Hatfield memorial lectures, Volume 2

The Hatfield Memorial Lectures Volume II

Dr William Herbert Hatfield FRS, 1882-1943. An informal picture taken in 1942 at Brincliffe House, Sheffield (from the Firth Brown Photographic Collection, reproduced by permission of the Kelham Island Museum).

The Hatfield Memorial Lectures Volume II Edited and Foreword by

Peter Beeley

MANEY publishing

B0771 First published in 2001 by Maney Publishing for The Institute of Materials 1 Carlton House Terrace London SWl Y 5DB © The Institute of Materials 2001 All rights reserved Maney Publishing is the trading name of w. S. Maney & Son Ltd Hudson Road Leeds LS9 7DL ISBN 1 902653 51 3

Typeset in the UK by Dorwyn Ltd, Rowlands Castle, Hants Printed and bound in the UK at The University Press, Cambridge

Contents Foreword Printed Sources

VII IX

Introduction Observations on William Herbert Hatfield FRS, A. G. Quarrell, 1963

1 3

Metallography and the Structure of Iron and Steel The Decomposition of Austenite by Nucleation and Growth Processes, R. F. MeW, 1948 Trends in Metallurgical Research in the United States, E. C. Bain, 1955 The Mechanism of Formation of Banded Structures, P. G. Bastien, 1957 Phenomena Occurring in the Quenching and Tempering of Steels, G. V. Kurdjumov, 1959 Metallography - A Hundred Years after Sorby, A. G. Quarrell, 1963 Intermetallic Chemistry of Iron, W. Hume-Rothery, 1965 The Status of the Metallurgy of Cast Irons, H. Morrogh, 1967 Metallic Chemistry in One, Two and Three Dimensions, L. S. Darken, 1969 The Heterogeneity ofSteel,J. W. Menter, 1970 Ferrite, R. W. K. Honeycombe, 1979 Clean Steel, Dirty Steel, J. Nutting, 1988

5

247 283 315 343

Author Index Subject Index

363 365

V

7 45 95 117 163 193 223

Foreword

The background to the publication of selections of the Hatfield Memorial Lectures in book form was outlined in the foreword to the first volume, published by the Institute of Materials in 2000 and devoted to the themes of properties, behaviour and applications of materials. This second volume is based on a single theme, containing lectures primarily concerned with aspects of metallography and especially with the structure of iron and steel. William Herbert Hatfield was born in 1883 and died in 1943 after long and distinguished services to metallurgy. The annual Lecture was instituted in 1944 and the series, beginning in 1946, continues successfully today. Biographical details of Hatfield were included in the foreword and introductory section to Volume 1, the latter consisting of the first lecture itself, presented by Dr George B . Waterhouse and devoted to Hatfield's own contributions to metallurgy. On a similar principle, the prologue to Professor Arthur Quarrell's Hatfield Memorial Lecture, the 15th, presented in 1963, has been brought forward to form the short Introduction to the present volume. Quarrell's first hand knowledge of Hatfield, through active participation in his important research committee work, provides a further valuable insight into the personality and professional life of this remarkable man. The main body of Professor Quarrell's lecture, 'Metallurgy a hundred years after Sorby', appears in sequence later in the book. The lectures in the present volume are arranged in a similar principle to that used in Volume I, namely in date sequence within the single group, so as to give a perspective of developments in the particular subject, and of the interest in it, at successive stages over the period since 1946. A further volume will be devoted to the themes of process metallurgy, research, and economic aspects of the iron and steel industry. In a few cases the lectures were not published but derived articles or reviews have been included instead. All the published sources are listed separately on page ix. Peter Beeley

P. R. Beeley DMet is a Life Fellow and former Senior Lecturer in Metallurgy in the University of Leeds VII

Printed Sources Listed below are the lecture numbers, titles and authors of each of the papers appearing in this volume. The original place and date of publication is also given. Third Lecture: The Decomposition of Austenite by Nucleation and Growth Processes, by Robert Franklin MeW PhD, DSc, DEng ]. Iron Steel Inst., June 1948, 113 Eighth Lecture: Trends in Metallurgical Research in the United States, by Edgar C. Bain ]. Iron Steel Inst., November 1955, 193 Tenth Lecture: The Mechanism of Formation of Banded Structures, by Paul G. Bastien, DSc ]. Iron Steel Inst., December 1957, 281 Twelfth Lecture: Phenomena Occurring in the Quenching and Tempering of Steels, by G. V. Kurdjumov J. Iron Steel Inst., May 1960, 26 Fifteenth Lecture: Metallography - a Hundred Years after Sorby, by A. G. Quarrell ]. Iron Steel Inst., July 1963, 563 Seventeenth Lecture: Intermetallic Chemistry of Iron, by W. Hume-Rothery ]. Iron Steel Inst., December 1965, 1181 Eighteenth Lecture: The Status of the Metallurgy of Cast Irons, by Henton Morrogh, FIM, FRS ]. Iron Steel Inst., January 1968, 1 Twentieth Lecture: Metallic Chemistry in One, Two and Three Dimensions, by L. S. Darken ]. Iron Steel Inst., January 1970, 1 Twenty-first Lecture: The Heterogeneity of Steel, by J. W. Menter, FRS ]. Iron Steel Inst., April 1971, 249 Twenty-ninth Lecture: Ferrite, by R. W. K. Honeycombe PhD, DSc, FIM Met. Sci., June 1980, 201 Thirty-sixth Lecture: Clean Steel, Dirty Steel, by J. Nutting Ironmaking Steelmaking, 1989, 16, (4), 219

IX

Introduction

Observations on William Herbert Hatfield FRS A. G. Quarrell

Prologue by Professor A.G. Quarrell to his Hatfield Memorial Lecture in 1963. This lecture, entitled 'Metallography - a hundred years after Sorby' is included in this volume, on p. 163.

Those who have delivered the Hatfield Memorial Lecture before me have rightly said that it is an honour to be invited to give it. It is a special privilege for me, for William Herbert Hatfield, or 'Billy Hatfield' as he was known far beyond his immediate circle of friends, was indirectly responsible formy coming to Sheffield and turning to metallurgy. As Chairman of the Alloy Steels Research Committee he authorised the creation of the temporary research assistantship that brought me here in 1936. Then I am conscious, as he was, of his close association with the University, and particularly with the Department of Metallurgy. He obtained his Associateship in Metallurgy in 1902, and the degrees of Bachelor, Master and Doctor of Metallurgy in 1908, 1912 and 1913 respectively. With his Associateship he gained the highly prized Mappin Medal - the hallmark of an outstanding metallurgist. I may well be the last person to give the Hatfield Memorial Lecture who also had the privilege of serving on research committees under Hatfield's chairmanship. On the Ingots Committee, the Alloy Steels Committee and the Hairline Cracks Sub-Committee, I had the opportunity of seeing the skilful way in which he handled them, and persuaded their members to cooperate in research. He did much to overcome secrecy and mistrust wherever it existed, and to develop such a spirit of cooperation within the steel industry that the British Iron and Steel Research Association was able to grow naturally from the research activities of the Iron and Steel Industrial Research Council, to which he had contributed so much. A man of great personality, Hatfield was likely to dominate any metallurgical gathering. There was something of the showman in his make-up, well illustrated by the way in which he resorted to carriage and pair when petrol rationing became severe during the second world war. But it would be wrong to attach too much importance to this side of his character. He made a considerable contribution to ferrous metallurgy, and for many years under his leadership the Brown-Firth Research Laboratories were unrivalled in this

3

4

Hatfield Memorial Lectures VoL II

country. He was quick to see the potentialities of new techniques and of new ideas, and was properly curious about the behaviour of metals. He was responsible for developing many special alloy steels and for improving them systematically over the years. Mechanical and physical properties, including some that had no apparent practical application at the time, were determined as a matter of course. It was largely as a result of this that when Whittle came forward with designs for his jet engine, steels of known and adequate properties were already available to meet his immediate demands, so making possible the first jet engine. I can speak from experience of Hatfield's kindness to younger scientists, and I recall how he stimulated them by deliberately asking provocative questions that would stretch their imaginations and their understanding. It was in character that the prize that he provided for the Sheffield Metallurgical Association should be intended to encourage its younger members. Tonight it seems specially appropriate that he should have named this award 'The Sorby Memorial Prize' . We may speculate on his reasons for choosing this title: he would have known that Henry Clifton Sorby was a descendant of the first Master Cutler and the grandson of another; that modem scientific metallurgy dates back to Sorby's pioneering work which had established the basis of metallography; he would have remembered that Sorby was a Vice-President of the University College of Sheffield when as a young student he, Hatfield, had received his Associateship in 1902; and he would probably have recalled Sorby's speech of welcome in 1905 to The Iron and Steel Institute, holding its Autumn General Meeting in the University of Sheffield, which had just received its charter. No doubt he was influenced by all these facts and wished to pay tribute to one of the greatest scientists Sheffield has produced. Certainly, we can be sure that he had a very high regard for Sorby, and so I feel it suitable to take as my title for the Fifteenth Hatfield Memorial Lecture, 'Metallography - a hundred years after Sorby'. I shall try to link it with the Sorby Centenary Celebrations to be held in Sheffield tomorrow by following a brief summary ofSorby's work with a review of modem metallography and some of its achievements.

Metallography and the Structure of Iron and Steel

THE

THIRD

HATFIELD

MEMORIAL

LECTURE

The Decomposition of Austenite by Nucleation and Growth Processes Robert Franklin Mehl, PhD, DSc, DEng. At the time the lecture was given Dr Mehl was Director of the Metals Research Laboratory, and Professor of Metallurgical Engineering, at the Carnegie Institute of Technology, Pittsburg, PA, USA. The lecture was presented in the Lecture Hall of the Royal Institution of Chartered Surveyors, 12, Great George Street, London, SW1 on 5th May 1948.

In considering a subject for this Hatfield Memorial Lecture, I have not been unmindful of what appeared to be the predilection of the Lecture Committee, that I should speak upon transformations in steel. This subject seemed appropriate to me, for the catholicity of Dr Hatfield's activities in the metallurgical field led him to manifest a deep and discerning interest in such matters, as the pages of the transactions of this Institute abundantly show, and the subject is thus one that would have interested him. Perhaps in this way this lecture may be taken as a tribute to him, to his long labours in the advancement of metallurgy, and to the beneficent influence he exerted. This subject is indeed one which has been pursued at some length in the laboratory with which I am associated, and I shall enjoy putting our views before you.

It is now quite well recognised that the properties of annealed and normalised steels are determined by the structures fonned at high temperatures, and it has been learned, increasingly in recent years, that these structures are determined by the rates of fonnation of the ultimate structures in the temperature ranges in which reaction occurs, and by the interdependent morphology. In the very important group of steels that are quenched for hardening, however, the success of this operation largely rests upon the ability to choose steels for which commercial quenching operations are rapid enough to avoid appreciable formation of the high temperature reaction products. Thus, obversely, the proper study of hardenability is in large part also the study of the rates of these high temperature reactions. These reactions consist of the formation of the pro-eutectoid constituents ferrite and carbide, the formation of pearlite and at lower temperatures the formation of bainite. They proceed by nucleation and growth, certainly for the first named constituents, and possibly also for bainite. It is to this subject, these nucleation and growth processes, that this lecture is addressed. The subject is not yet far advanced, but enough has been done to permit this appraisal and to anticipate those lines along which research might most profitably follow.

7

8

Hatfield Memorial Lectures VoL II

In considering the fundamental question of the kinetics of these reactions, detailed attention must be given to the morphology of the resultant constituents. In reactions within solids, kinetics and morphology are interdependent, for the detailed characteristics of kinetics determine the shape and spatial distribution of the products, and the latter in turn affect the overall rate of the reaction. Both rate and morphology are best studied by isothermal reaction, and this is the approach that will be used. The correlation of isothermal rates of reaction with the phenomena occurring on cooling is not a subject of extraordinary difficulty, and in some respects is well advanced.

ISOTHERMAL REACTION - THE TYPES OF STEEL At a temperature below the Ael temperature, austenite will decompose on cooling or may be decomposed isothermally. The isothermal rate, usually measured microscopically or dilatometrically, is represented by an isothermal reaction curve, in which the fraction transformed is plotted against time (Fig. 1). These curves, taken at a series of temperatures, are usually assembled into an isothermal transformation diagram, by the device of plotting the beginning of the reaction at, say, 1% reaction and the end of the reaction at, say, 99% reaction.

~

100

(I

eft 80

/

5

~bO ~ 40 ~ 20

I

I

./''1 10

TIME. S

Fig. 1

..•

102

A typical isothermal reaction curve.

Figure 2 shows such an isothermal transformation diagram for a eutectoid carbon steel. Above the knee of the curve at about 570°C pearlite forms, and below the knee, bainite. At a low temperature, martensite forms, represented on this diagram by a horizontal line to indicate the beginning of the reaction. The martensite reaction occurs substantially only on cooling, not isothermally; it does not form by nucleation and growth, but by shear unaccompanied by diffusion, and is not considered further in this lecture. Figure 3 is a similar curve for a hypo-eutectoid carbon steel, showing the formation of pro-eutectoid ferrite preceding the formation of pearlite, the amount of the pro-eutectoid constituent decreasing as the temperature decreases and approaching zero at the knee. Hypereutectoid steels have a similar type of isothermal diagram, with pro-eutectoid cementite replacing pro-eutectoid ferrite.

The Decomposition of Austenite by Nucleation and Growth Processes

o~--~--~--~----~ 10 J02 lOS 10

4

TIME S

°1

lOS

Fig. 2 Isothermal transformation diagram for a eutectoid steel with 0.79% carbon, 0.76% manganese, grain size 6 (US Steel Corporation Research Laboratories).

10

102 TIME.S

f03

104

9

IQ5

Fig. 3 Isothermal transformation diagram for a hypo-eutectoid steel with 0.50% carbon, 0.91% manganese, grain size 7-8 (US Steel Corporation Research Laboratories).

The isothermal decomposition of alloy steels may be represented in a similar fashion. In Fig. 4, which shows an isothermal diagram for a 4% nickel steel, it is observed that nickel merely displaces the curve to the right, in terms of cooling, providing deeper hardening, while the shape of the curve is not changed. Figure 5 shows an isothermal diagram for a chromium-nickel-molybdenum steel; here the pearlite region is seen to be displaced to the right, but also the curve .now is more complex, with a pro-eutectoid constituent resembling ferrite forming down to quite low temperatures. This complexity in the form of isothermal diagrams for alloy steels reflects the complexity in the mechanism of the reactions which such steels exhibit.

~

~"400 ~

~

~M.:.;;;s--l~--+~

~lOO·~--~--~--~--~~

•... &.oLJ

°1

10

102

TIME,S

103

105

Fig. 4 Isothermal transformation diagram for a steel with 0.55% carbon, 0.33% manganese, 3.88% nickel, grain size 8-10 (US Steel Corporation Research Laboratories).

°1

to

Fig. 5 Isothermal transformation diagram for a steel with 0.42% carbon, 0.76% manganese, 1.79% nickel, 0.80% chromium, 0.033% molybdenum, grain size 7-8 (US Steel Corporation Research Laboratories).

Hatfield Memorial Lectures Vol. II

10

These are now elementary matters, but they are introduced here because the reaction characteristics of each of these main types of steel will be considered in some detail.

THE CHARACTERISTICS OF PROCESSES OF NUCLEATION AND GROWTH The decomposition of austenite at sub equilibrium temperatures, either on cooling or isothermally, is a typical heterogeneous reaction, proceeding by the formation of nuclei and their subsequent growth. All heterogeneous reactions, for all states of aggregation, proceed in this way, except those solid-solid reactions which, like martensite, proceed by shear. In the most general case, as shown in Fig. 6, the nuclei form at random and each grows. Nuclei continue to form with time, and each grows to a nodule, striving to form a spherulite, with growth ultimately restricted by impingement with other growing nodules.

Fig. 6

Diagram depicting the progress of transformation in a nucleation and growth process. Four stages at equal intervals.

This ideal case of nucleation and growth, when the rate of nucleation and the rate of growth are both constant with time, may be expressed mathematically as:

ji( t ) -- 1 - e - !!3 NG3t

4

where the fraction transformed f(t) is given in terms of the rate of nucleation N, the rate of growth G and the time t. Such a reaction equation produces a reaction curve of the type shown in Fig. 7, identical in character with the experimental isothermal reaction curve for pearlite shown earlier (Fig. 1). This treatment applies well, indeed nearly perfectly, to pearlite formed at high temperatures. In other cases, modifications must be made in the reaction rate equation, but these modifications relate only to alterations in the analytical form, including allowance for the formation of reaction products which have geometrical forms other than that of

The Decomposition of Austenite by Nucleation and Growth Processes 81·0 '-

N -1()()()Ieu em

C-3xO)cm/s

/

./

11

Is /..-

/

/

~

V

200 400 TIME, S Fig. 7

Calculated reaction curve for general nucleation.

spherulites, and do not affect the basic principle, that the rate of nucleation and the rate of growth determine the isothermal rate of reaction. The factors that determine the rate of nucleation and the rate of growth, and the contribution that theory can make with respect to each of these rates, must therefore be studied.

THE FORMATION OF PEARLITE IN EUTECTOID

CARBON STEELS

The formation of pearlite in eutectoid carbon steels is a good point of departure; as compared to hypo-eutectoid or hypereutectoid steels, pearlite is relatively simple, and it illustrates many of the basic principles for all compositions. Consider first the formation of pearlite near Ae1• The isothermal reaction curve is that shown in Fig. 1. Microscopically it is observed that nuclei form at the austenite grain boundary and grow to large nodules. Each of these nodules is composed of the pearlite colonies of Belaiew (Fig. 8), areas formed as a unit, usually with but one direction of lamellar, in which the ferrite and the cementite have each a single orientation. The original nucleus, a cementite platelet, appears and in time thickens, then a parallel lamella of ferrite deposits and in turn thickens, then cementite again, and so on, all lamellae meanwhile also growing edgewise, creating the pearlite colony. The formation of a pearlite colony thus seems to proceed by edgewise growth, that is, in the direction of the lamellae, and also by sidewise nucleation and growth. This first colony in turn nucleates other colonies, so that a nodule forms by the successive formation of colonies (Fig. 9). The rateat which these nuclei form, and the rate at which they grow have surrendered to measurement.

It has been shown that the number

of nuclei formed in a given period of

time for a steel in different grain sizes varies with the extent of the grain boundary surface, i.e. that the basic value of the rate of nucleation is essentially the rate per unit grain boundary area; except for substantial austenite heterogeneity, the rate of nucleation

12

Fig.8

Hatfield Memorial Lectures Vol. II

Pearlite colonies in a eutectoid steel transformed at 715°C. Electrolytically polished and etched with Vilella's reagent x2500.

(I.) Initial

Fe,( nucleus (2) Fe,( plate fulf-qrow~ acFe so« nucleated

(3)0< Fe plate now full-qrown new Fe,C plates nucleated

(-4) New FelC rucleus of different orientatiOn forms at surface of colony durinq s~~ise oocleation (5) New colony at advanced ~~~~r~?nqtq~~~~ staqe of qrowth

Fig. 9

Nucleation and growth of pearlite colonies.

within the grain is of no consequence - indeed this may be employed as a quantitative test of austenite heterogeneity; inclusions, though active occasionally as inoculants, in fact play an exceedingly minor part. It is, of course, in this predominating exclusiveness of grain boundary nucleation that the effect of grain size on reaction rate and the correlative hardenability lies. On this basis the effect of grain size on hardenability has in fact been appraised by Grossmann. The rate of nucleation, surprisingly, is not a constant with time, as Fig. 10 shows. This has been observed in all cases in the decomposition of austenite, even for heterogeneous austenite, and also in the somewhat similar case of nucleation and growth during recrystallisation; it may, in fact, be a general phenomenon in nucleation processes (investigators of the more conventional types of nucleation processes appear not to have investigated the point). The rate of growth of a pearlite nodule as measured is the rate of growth of successive colonies. Measurements show this to be very close to the true rate of edgewise growth

The Decomposition of Austenite by Nucleation and Growth Processes 1400

13

oV

1200

I

000

~

800 600

if

I{

/

400 200

o Fig. 10

,."

2

~~

~ b

TIME.S

10

14

Rate of nucleation of pearlite as a function of time. Eutectoid steel reacted at 680°C.

(Fig. 11) . Now pearlite represents the complete decomposition product of austenite, that is, the composition of austenite away from the pearlite-austenite interface does not change; decomposition is initiated and completed at the moving pearlite interface. Accordingly, it is not surprising that the rate of growth of pearlite is constant with time. Moreover, pearlite grows across austenite grain boundaries with no perceptible retardation, and is in no way affected by austenite heterogeneity. In a word, it is a fundamental characteristic of the rate of growth that it is not structure sensitive. The rate of nucleation, however, far higher in the disorder of the grain boundary than within the grain, and strongly accelerated by austenite heterogeneity, is sharply structure sensitive.

Ferrite

Low carbon concentration ::ZZZ:Zz:z:z:z:l:ll:Z:ZZZZ=~i-in austenite -C,

Cementite ~:zz:zzzz:z:z:z:z:zz:zz:_-4-Hiqh carbon concentration in austenite· (2

Austenite

Fig. 11

Simplified model showing the edgewise growth of pearlite.

N ear the Ael temperature the rate of nucleation is small compared to the rate of growth, with the result that the pearlite nodule grows very large (Fig. 12) absorbing many austenite grains, before it impinges upon another growing nodule. Even though the nuclei originate at grain boundaries, their distribution is on such a fine scale with respect to the final nodule size that nucleation may be considered as effectively random. Taking an average rate of nucleation to represent the rate changing with time, the reaction at high temperatures thus very closely approximates the idealised case first pictured, and the isothermal reaction rate curve may be closely calculated from the rate of nucleation per unit austenite grain boundary surface and the rate of growth.

14

Hatfield Memorial Lectures Vol. II

Fig. 12

Pearlite nodules formed in eutectoid steels after 9 min at 680°C. Austenite grain size shown by inset sketch xmo.

At temperatures near the knee of the isothermal diagram, the morphology of pearlite changes. Microscopic evidence (Fig. 13) gives direct proof that nucleation is predominantly at the grain boundary; here again, and in fact throughout the temperature range in which pearlite forms, the rate of nucleation must be expressed as the number of nuclei formed in unit time per unit austenite grain boundary surface. Near the knee of the isothermal diagram the rate of nucleation per unit austenite grain boundary surface is very high, with many pearlite nodules forming per grain. The result of this is that the mode of growth is quite different: each nodule grows but a small way before it impinges upon another growing nodule, whence these nodules grow toward the centre of the grain roughly in the shape of sectors. It is clear that the morphology of pearlite nodules is dependent upon the ratio of the rate of nucleation and the rate of growth; this ratio is high at temperatures near the knee, and low at temperatures near Ae1• The morphological characteristics of the reaction at temperatures near the knee distinguish the reaction from the ideal assumption of random nucleation and random

Fig. 13

Grain boundary transformation in eutectoid steel upon quenching from 800°C xl00.

The Decomposition

of Austenite; by Nucleation and Growth Processes

15

impingement originally assumed in deriving the reaction equation; but the reaction equation can be modified to take into account this different type of growth geometry, and this has been done. Whereas in truly random nucleation and growth the form of the isothermal reaction curve is always of the same shape, in the case of reaction near the knee a series of shapes may eventuate, all predictable, depending upon the relative values of the grain size, the rate of nucleation and the rate of growth. No mystery on this aspect of the formation of pearlite remains.

The Rate of Nucleation - Experimental Evidence and Theory Rates of nucleation have been measured in a number of pure and commercial ironcarbon alloys of eutectoid composition, over the full temperature range. One of the curves obtained is shown in Fig. 14 and represents an average rate of nucleation, suppressing the variation of this quantity with time. All curves thus far measured are similar to this. It will be observed that the rate is presumably zero at the equilibrium temperature and reaches a high value near the knee of the isothermal diagram.

550~'-;--~ra~..;---L..rO';""'"3--~--J....----'Ol 0-4

Fig. 14

RATE OFGROWTH MM1S 10-2 100 JQ2 RA1E OF NUCLEATION. NUCLEI /Cu. MM/S

J04

Rate of nucleation and rate of growth in a eutectoid steel (0.78% carbon, 0.63% manganese, ASTM grain size no. 5.25) as a function of reaction temperature.

The active nucleus in this reaction is almost certainly cementite. The evidence for this is of several kinds. Determinations of the orientation of ferrite in pearlite with respect to the parent austenite show it to be quite different from the orientation when ferrite is known to nucleate directly from austenite, as in the 'Y-u transformation in pure iron or as in the formation of pro-eutectoid ferrite from austenite; the only conceivable explanation is that cementite acts as the nucleus and that the ferrite, forming subsequently, derives its orientation from the orientation of the cementite. Moreover, it is observed that residual cementite in austenite is a powerful inoculant for pearlite, whereas ferrite is not. It is in fact observed that the cementite in pearlite in hypereutectoid steels is continuous with the grain boundary pro-eutectoid cementite, whereas the ferrite in pearlite is not continuous with pro-eutectoid ferrite in hypoeutectoid steels. The variation of the rate of nucleation with temperature is in accord with formal nucleation theory. This theory, developed originally by Gibbs to apply to heterogeneous

16

Hatfield Memorial Lectures VoL II

reactions generally, provides an expression for the rate of nucleation and a formal treatment of its variation with temperature. Reactions proceed only when the free energy of the system as a whole decreases; it is not necessary that the free energy go to the lowest possible, that is, the equilibrium value. This is a highly important qualification for metallurgists, for it provides that intermediate stages may form, characterised by values of free energy change less than the maximum; chemists will remember Ostwald's law of stages; and in this circumstance lies the great opportunity, which metal systems fully exercise, of forming a multitude of intermediate stages. The cementite in pearlite itself is unstable, and forms only because its rate of formation is greater than that of the stable phase graphite; martensite is not a stable phase, but a transition structure having many features in common with the transition lattices in age hardening systems responsible for the hardening that occurs; and any disperse system, such as pearlite, is unstable with respect to a system of the same phases in which each constituent is a single crystal - in pearlite appreciable amounts of energy reside in the ferrite-cementite interface; when pearlite is spheroidised it is this energy which is the driving force. When a new phase forms, the free energy of the system decreases, that is, free energy is made available, the amount being determined by the volume of the new phase formed. The creation of the new phase involves the formation of an interface between the new phase and the parent phase, and this requires the expenditure of energy, furnished from the free energy made available. If the particle of the new phase is very small, the amount of free energy available is small and inadequate to fonn the interface. The growth of such a phase would require an increase in the free energy of the system, which is impossible; such nuclei thus cannot form and grow. Since the volume and thus the free energy increases with the cube of the radius while the interface increases with the square of the radius, there will be a given radius, a larger radius, at which the free energy made available by this larger volume exactly equals the energy requirement for the formation of the interface. Such a nucleus is stable, and it can grow, as can any larger nucleus. The mechanism by which these nuclei come into being, associated with normal fluctuations in concentration and the corresponding diffusion rates and interaction energies, is not well understood. The stable nucleus size should, according to this scheme, decrease with temperature, and the rate of nucleation calculated from it should increase with decrease in temperature. These considerations, relatively easy to apply to simple heterogeneous reactions, such as the condensation of vapour, encounter major and still unresolved difficulties in solidsolid reactions. This is primarily though not exclusively because the values of the interface energies between solids are entirely unknown and thus far unmeasurable. Since one solid phase forming from another is always oriented, that is, the crystal faces in juxtaposition are selected faces, the energies of specific interface couples would have to be treated, and since lattice coherency between the two phases usually obtains, strain energies must modify the interface energy. In a word, thus far only a thoroughly qualitative theory with respect to the rate of nucleation has been developed. When considering alloy steels the difficulties in this respect will increase. The study of the rate of nucleation is now restricted to a purely inductive approach; this is regrettable,

The Decomposition

of Austenite by Nucleation and Growth Processes

17

for the rate of nucleation is the more important variable in determining the rate of decomposition of austenite, and the host of engineering and practical matters depending upon that.

The Rate of Growth The lesser variable, the rate of growth, may be much more readily treated in a quantitative way. Once the growth of a colony is established, by- the mechanism postulated earlier, further growth is edgewise, a phenomenon clearly apparent in the radial growth near the knee of the isothermal diagram, and quite obvious where pearlite nodules are formed partially at one temperature and then grown farther at another. The rate of growth is easy to measure, and measurements are available on a number of commercial and pure iron-carbon alloys. The rate of growth increases with decreasing temperature of reaction, though not so rapidly as the rate of nucleation. One of the curves obtained is shown in Fig. 14; this figure serves to compare the values for the rate of nucleation and the rate of growth. Consider next the process of edgewise growth. Growing edgewise, that is, in the direction of the lamellse, as shown in Fig. 11, it is obviously a process of the segregation of the carbon from homogeneous austenite into carbide plates. Evidently there must be carbon concentrations set up in the neighbourhood of the moving interface, with high carbon concentrations in front of the ferrite plate and low carbon concentrations in front of the cementite plate, providing the necessary downhill gradient for the diffusion of carbon. This gradient must be known if the rate of growth is to be calculated. Brandt has recently solved this problem mathematically, apparently in a thoroughly satisfactory manner, assuming that the rate of diffusion of carbon in austenite does not vary with concentration. The concentration extremes present a difficult problem. The problem is usually compared to that of the growth of a crystal from a supersaturated solution, in which case it has generally been assumed that the solution at the crystal surface has the equilibrium value of saturation; this may require modification, for, strictly interpreted, under such circumstances no growth could occur, but the concentration at the point of growth may be very near to the saturation concentration. This argument was applied years ago to the formation of eutectics, and in recent years Hultgren has stated it clearly for eutectoids. Hultgren extrapolated the equilibrium curves below the Ac1 temperature, as shown in Fig. 15, and stated that this gives the limiting concentrations desired, that is, the concentration in austenite in front of the ferrite is that given by the extrapolation of the A3 (GS) curve, while that in front of the cementite is given by the extrapolation of the ACID curve. This is shown in Fig. 15 for a given reaction temperature; the concentration in front of the ferrite lamellae at the temperature chosen is c2, and that in front of the cementite lamellae is c1. As already suggested, the limiting concentrations may merely approach these extrapolated values as limits. Brandt assumed values of the limiting concentrations as

18

Hatfield Memorial Lectures VoL II

bOO

Fig. 15

Ferrite + cementite

Hultgren's extrapolation of GS and ES curves for iron-carbon alloys.

extrapolated in calculating the rate of growth. The difference between these concentrations, and half the distance between the lamellse, the interlamellar spacing, define the concentration gradient. This gradient and the diffusion coefficient together determine how fast carbon is delivered to the growing carbide plate and thus determine the rate of growth. On decreasing temperature of reaction, the limiting concentrations increase, as the extrapolation shows, and the interlamellar distance decreases, both contributing to an increased rate of growth; the diffusion coefficient decreases also but not sufficiently to compensate for the first two factors, and thus the rate of growth increases as the temperature decreases. Brandt's fundamental work may be modified by taking into account the variation of the diffusion coefficient with carbon concentration. Recently, Batz and Wells, in the author's laboratories, have extended earlier measurements on the diffusion coefficient of carbon in yiron to high carbon concentrations, as shown in Fig. 16. These measurements show that the diffusion coefficient is increasing rapidly as the concentration increases. If this curve is extrapolated to the extreme limiting concentrations required by the Hultgren extrapolation, it reaches high values. Introducing these values into the Brandt solution, (Fig. 17) the disagreement he found at low temperatures, amounting to a lower rate than measured, is removed and the agreement is quite good. The rate of growth of pearlite in carbon steels is thus well understood. The interlamellar spacing itself presents a basic problem. Careful measurements of the spacing as a function of temperature provide curves of the type shown in Fig. 18. The interlamellar spacing decreases as the temperature of formation decreases. It is difficult to say whether this spacing decreases linearly with temperature, or in some other way, for example, exponentially. The basic problem is: What determines the spacing? Until recently no persuasive or even attractive argument on this point has existed; however, it has

The Decomposition

of Austenite by Nucleation

and Growth Processes

19

v.-

s 4

3 2~~

a.

~

~

V

Vi

J

4 5 b 2 I 3 CARBON CQ.JCENTRATION ATOMIC

60010-4

10-3 VELOCI1Y

0/0

Fig. 17

Fig. 16 Diffusion coefficient of carbon in austenite as a function of carbon concentration (temperature 1195°C ± SOC).

OF GROWTH,

10-2

10-1

MM/S

Comparison of observed and calculated rates of growth of pearlite.

recently been proposed by Zener, in a very interesting paper, that the spacing is set by the free energy available in the reaction. The area of the ferrite-cementite interface in unit mass of pearlite obviously increases with decrease in the spacing; the lower the temperature of formation the greater the free energy available, and this is the energy that must be employed in creating the interface; accordingly more energy is available for this purpose at lower temperatures and the spacing can thus decrease. What it is that determines how much of the free energy is allotted to this purpose, and exactly what the temperature variation of the spacing should thus be, is not wholly clear. It might be observed in passing, that although thermodynamic measurements are extremely difficult in these reactions, owing to the smallness of the heat energies involved, precise data would be extraordinarily useful; they would remove such arguments comfortingly from the depths of speculation.

-,

lOxleY

8x

0

10

Id Id

"

"

<, ~

0

"~

Id 80~

x10-2 4x 0-2

Fig. 18

x ~

.r

2x 10-2

IxIO-2

0

Interlamellar spacing of pearlite as a function of the reciprocal of the undercooling.

20

Hatfield Memorial Lectures VoL II THE FORMATION OF PEARLITE IN EUTECTOID ALLOY STEELS

Alloying elements exert a complicating effect upon the pearlite reaction, leading ultimately to phenomena that are difficult to understand, even qualitatively. As to nucleation in alloy steels, it will again be assumed, and there is evidence for it, that the carbide phase, whatever it may be, nucleates pearlite. The addition of an alloying element to austenite creates a smaller probability of the chance association of atoms of the right composition to form a nucleus and, perhaps of more importance, alters the energy relationships between austenite and the nucleus. A smaller free energy change, other factors equal, would give a lower rate of nucleation, and any special requirement of composition of the nucleus would give the same effect. Inasmuch as the carbide in alloy pearlite is usually modified cementite, or some quite different carbide, the original nucleus presumably also departs in composition from that of ordinary cementite. Moreover, since nucleation in this case also is nearly exclusively at the grain boundary, these circumstances must be considered as they apply at the grain boundary. Since the days of Gibbs, it has been known that solutes must be concentrated at boundaries if this provides lower surface energies, and conversely. While probability calculations of atom distribution might be made, little can be done concerning the question of energy changes and of surface concentrations; any argument can be only qualitative and the subject must be approached from experiment. With the exception of cobalt, all alloying elements when in solution in austenite appear to decrease the rate of nucleation of pearlite. This conclusion is drawn from observations that are qualitative (though unmistakeable), for apart from cobalt no accurate determinations have been made; the work on eutectoid molybdenum steels by Blanchard, Parke and Herzig, shows that molybdenum apparently decreases the rate of nucleation by a factor perhaps as large as 1000. The marked effect of chromium in decreasing the rate of formation of pearlite, especially at the lower chromium concentrations where on a percentage basis the effect is the greater, is a parallel case. Measurements on molybdenum steels will shortly be available. The case of cobalt steels is especially interesting. Hawkes made careful measurements on these steels, showing that the rate of nucleation of pearlite is appreciably increased by cobalt. Perhaps the proposal referred to later in considering growth may apply here also, that cobalt increases the free energy change of the pearlite reaction, thus making more energy available for nucleation and thus increasing the rate; if the remaining alloying elements should have the opposite effect, their retarding effect upon nucleation would be at least in part explicable. When the amount of the alloying element is high, a new set of phenomena appears. For example, at molybdenum contents above 0.50%, the carbide in the pearlite is no longer orthorhombic cementite, but is instead a face centred iron-molybdenum carbide. Similar behaviour is exhibited by high chromium alloy steels, as Lyman and Troiano have shown. Although these conclusions are drawn from studies of samples that have reacted to form substantial amounts of pearlite, it may, though perhaps not with utter certainty,

The Decomposition

of Austenite by Nucleation and Growth Processes

21

be assumed that the original nucleus is also a modified carbide. The change to a different carbide must, of course, mean that the rate of nucleation of this new carbide exceeds that of cementite. If the composition of the nucleus could be predicted from the ternary equilibrium diagram the problem would be greatly simplified, but as will be seen when the growth of pearlite is considered, this is most certainly not so; the problem is one of kinetics, not equilibrium, as indeed it is throughout this whole subj ect. The effect of alloying elements upon the rate of growth of pearlite is known better, and may be treated more nearly quantitatively than the effect upon the rate of nucleation; at least, the problem can be more rigorously formulated. Again except for cobalt, all alloying elements decrease the rate of growth of pearlite; many qualitative observations show this and direct measurements on manganese and molybdenum steels and semi-quantitative measurements on others have been made. In carbon eutectoid steels, reacted at 680°C, variations in manganese from 0.2 to 0.8% cause the rate of growth to decrease to one-fifth; new work by Parcel on molybdenum eutectoid steels reacted at 650°C, shows the rate of growth to decrease to one one-hundredth as molybdenum increases from 0.02 to 0.52%. As already mentioned, the rate of growth should be determined by the interlamellar spacing, the concentration extremes and the diffusion coefficients of whatever elements diffuse. As to the diffusion coefficient, measurements are available that show that alloying elements have only a very small effect upon the rate of diffusion of carbon; but if the alloying element itself must diffuse in forming pearlite, then, owing to the relatively very slow rate of diffusion of the alloying element, there may be one reason for the slow rate of growth. It has been said that alloying elements cannot diffuse during the formation of pearlite, for, it is asked, why should pearlite which normally grows rapidly wait for the segregation of the slowly diffusing alloying element? The question should be answered by experiment. The available evidence, for example, in partition studies in molybdenum steels, shows that the cementite in low molybdenum pearlite does indeed contain molybdenum. The molybdenum content of the cementite carbide in isothermally reacted steels is in fact greater than that which can be induced into the carbide after far longer tempering of quenched steels at the same temperature, which surely argues against immediate and thus rapid molybdenum segregation just behind the growing pearlite interface. There is, in addition, evidence that manganese diffuses to the cementite in forming pearlite in manganese steels. It is obvious that when the carbide in the pearlite formed is a special molybdenum carbide or special chromium carbide, the diffusion of these alloying elements must occur during growth. Even if it be accepted that the diffusion of alloying elements retards the growth of pearlite, the effect of alloying elements on the interlamellar spacing, which, noted below, provides a like retardation, is also a contributing factor; the evidence is inadequate to appraise the relative contributions of these two factors. Little or nothing can be said concerning the concentration limits in the austenite in advance of the growing pearlite interface in alloy steels, that is, the extrapolation of the solubility curves in alloy steels is wholly conjectural. Brandt's calculations, however, suggest that the effect of conceivable changes in these limiting concentrations on the rate

22

Hatfield Memorial Lectures VoL II

of growth cannot be large, and it thus seems that exploration of this possibility would probably be fruitless. The interlamellar spacing, however, presents some interesting aspects. It will be remembered that a greater spacing means a lower rate of growth, other factors remaining unchanged. At the same temperature, nickel, manganese and molybdenum decrease the rate of growth, and correspondingly increase the interlamellar spacing, whereas cobalt, which increases the rate of growth, decreases the spacing; the decrease in spacing caused by cobalt accounts quantitatively for the increased rate of growth observed. The correlation of spacing with the rate of the pearlite reaction and presumably with the rate of growth, is close. Molybdenum has the greatest effect upon the rate of the pearlite reaction near the knee of the isothermal diagram where the effect of molybdenum on the spacing is large, and least near Ael where the spacing seems unaffected; as little as 0.15% of molybdenum has a detectable effect upon the interlamellar spacing. This correlation of the spacing and the rate of growth appears to be important. Zener argues that there is reason to believe that cobalt increases the free energy of the pearlite reaction as compared to that in carbon steels, and thus provides more energy for the creation of the ferritecementite interface in pearlite, resulting in a decreased spacing. On this basis nickel, manganese and molybdenum might be expected to decrease the free energy somewhat. If the comparable argument is employed in nucleation, as suggested above, nucleation and growth have some fundamental similarities. In carbon steels, pearlite forms down to the knee of the isothermal diagram, and no farther, joining there, imperceptibly in rate, with the formation of bainite. In alloy steels, however, pearlite may fonn a knee of its own, that is, its rate of formation passes through a maximum as the temperature decreases, as in the cases of molybdenum steels and chromium steels. The process, of course, is still one of nucleation and growth, but there have been no measurements of the rates in such a case, nor of interlamellar spacings, though they are under way. Various suggestions have been made concerning mechanisms that will provide for a maximum in the rate of nucleation and the rate of growth at some subcritical temperature, but they lack the virtue of submitting to test. In view of the complexities in alloy steels where the phenomenon is observed, it would appear preferable at the moment to pursue the subject experimentally until the problem itself is clarified in terms of known rates of nucleation and growth. That knowledge of equilibrium helps but little in such studies is nowhere shown more clearly than in the reactions of alloy steels. It has been observed in both molybdenum and chromium steels that the partition of the alloying element between ferrite and cementite in pearlite varies with the temperature at which the pearlite forms. Moreover, the partition coefficient when the pearlite reaction is just complete changes as this pearlite is held for long times at the same temperature, striving towards the equilibrium partition. Kinetics determine the initial partitions, not equilibrium, and of course it is this circumstance that makes the study both new and difficult. The complexities in alloy steels that form pearlite do not cease here, however. Hitherto, in considering eutectoid carbon steels, it has been decided, on evidence, that no

The Decomposition

of Austenite by Nucleation and Growth Processes

23

compositional changes occur in the austenite during the growth of pearlite, but Lymanand Troiano have recently shown that in the formation of pearlite in chromium steels, the unreacted austenite is depleted in carbon as the reaction proceeds. While the reaction is, of course, still one of nucleation and growth, evidently, in deriving kinetics, such reactions must be clearly distinguished from the more ordinary ones. It should be remarked finally, that these reactions do indeed proceed by nucleation and growth, whatever their complexities, and that an explanation of the isothermal reaction rates, which compose the isothermal diagram, must be sought in the factors which determine the rates of nucleation and growth which themselves completely fix the isothermal rate.

HYPO-EUTECTOID

STEELS

Hypo-eutectoid steels are more important, practically, than eutectoid steels, for they constitute the great bulk of all steels manufactured. Some of the principles advanced to account for the formation of pearlite may be used here, and to these must be added others that specifically apply to this type of reaction. The study of rates of nucleation and growth in hypo-eutectoid steels is still in its infancy, and must await a thorough understanding of the factors controlling morphology. It is to this latter subject that this section will be largely devoted. The elementary features of the rate of isothermal reaction of hypo-eutectoid steels are well known. Figure 3 portrays the rates of reaction as assembled in an isothermal diagram. Hypo-eutectoid steels first reject pro-eutectoid ferrite and, at temperatures below Ae1, later form pearlite. The amount of pro-eutectoid ferrite decreases with decrease in temperature; in medium carbon steels it decreases to zero or nearly zero at the knee of the curve, which is the equivalent of saying that pearlite formed at lower temperatures contains less carbon. The formation of pro-eutectoid ferrite proceeds by nucleation and growth, but growth leads to microstructures geometrically more complex than in eutectoid steels. The isothermal rate of formation of pearlite must be treated in terms of the rates of nucleation and growth, but owing to the complex morphology of pro-eutectoid ferrite, the isothermal reaction equation cannot be so simply derived as in the case of pearlite where simple spherulites of the product could legitimately be assumed. Accordingly, it is especially important in these steels to consider morphology in attacking the question of kinetics. Carpenter and Robertson in their splendid work of nearly twenty years ago studied morphology in samples reacted during cooling. The new work reported here relates to isothermal reaction, complementing and extending the work of Carpenter and Robertson; it is largely the work of Mr Arthur Dube, who has been associated with the author.

The Morphology ofPro-Eutectoid The morphology

of pro-eutectoid

Ferrite

ferrite is complex. It varies with temperature,

content, time and austenite grain size. Consider first the morphology

carbon

of a 0.32%

24

Hatfield Memorial Lectures VoL II

commercial carbon steel, reacted at different temperatures. The upper equilibrium temperature, Ae3, is about 795°C and the Ael temperature is about 720°C. The reaction just under the Ae3 temperature, as Carpenter and Robertson showed, is the formation of small ferrite crystals at the austenite grain boundaries. Figure 19 shows this steel, with an austenite grain size of ASTM No.7, that is, of a mean grain area of 0.001 mm-', reacted for 30 s at 725°C. In the subsequent discussion this steel will be designated as the fine grained steel. As time goes on the reaction consists simply in the formation of new ferrite crystals, and thickening and prolongation of the ferrite along the grain boundary. This simple morphology is obtained in the temperature range down to about 685°C. At 675°C the initial structure is again composed of small crystals at the grain boundary though the greater number of these in a given time demonstrates an increased rate of nucleation with decreasing temperature. As the reaction proceeds isothermally, these crystals thicken and extend along the grain boundaries until a point in time is reached when plates of ferrite are observed which grow from the grain boundary ferrite toward the centre of the grain (Fig. 20). These plates ultimately thicken somewhat and agglomerate, and the formation of pearlite sets in, so that the final structure shows only massive ferrite and pearlite. At 650°C the same sequence is observed, but agglomeration is minor, apparently interrupted by the formation of pearlite which is more rapid at this temperature, and the Widmanstattcn pattern, though not of high perfection, is maintained in the final microstructure. At lower temperatures the same sequence is observed, with less and less ferrite appearing, until a temperature of about 600°C is reached, when the amount of pro-eutectoid ferrite is so small that the reaction does not proceed beyond the grain boundary stage. In this same steel in the coarse grained condition, with an austenite grain size of ASTM No.1, that is, a mean grain area of 0.062 mm-', profound alterations in morphology appear. At 750°C the familiar grain boundary ferrite appears, outlining the grains (Fig. 21) without subsequent alteration. On reaction at 725°C the same initial stage obtains, but

Fig. 19

Ferrite crystals formed at grain boundaries after reaction at 725°C for 30 s in a hypo-eutectoid

Small grains xSOO.

steel.

The Decomposition of Austenite by Nucleation and Growth Processes

Fig. 20

25

Ferrite plates growing form rain boundary crystals in hypo-eutectoid steel after reaction at 675°C for 10 s. Small grains x1000.

this is followed by the formation of ferrite plates, which ultimately thicken. The Widmanstatten pattern in this steel thus appears at a higher temperature than in the fine grained steel. At 700°C the same features are shown, except that more ferrite appears, and ferrite grains have appeared within the austenite grain (Fig. 22). These grains are Widmanstatten in character and are characteristically more precisely Widmanstatten in form than the plates which grow from grain boundary ferrite. At lower temperatures, below Ae1, down to about 650°C, the same sequence is observed, and as is to be expected, less ferrite appears in the fully reacted steel, for again the formation of pearlite interrupts the growth of ferrite. The total amount of pro-eutectoid ferrite formed below Ael is less in coarse grained steel than in the fine grained steel similarly reacted, as was noted by Carpenter and Robertson on steels reacted during cooling. An explanation of this phenomenon will be given later. At 650°C the reaction to pearlite interrupts the ferrite reaction at such an early stage that only grain boundary ferrite forms. As the reaction temperature is lowered further, this effect increases until ferrite is no longer detectable. Morphological sequences have also been followed in coarse and fine grained steels with 0.48% and with 0.62% of carbon. In the higher carbon steels, Widmanstattcn figures do not appear because of the early interruption of the reaction by the formation of pearlite. These structural results for the three steels noted are plotted in Fig. 23, which represents the final microstructures of the isothermally reacted steels. It is clear that temperature, grain size and percentage of carbon are all variables which determine the final microstructure.

Theory of Morphology These morphological features may be in part understood. The far higher rate of nucleation at the grain boundary is similar in origin to that observed in pearlite, and the rapidly

26

Hatfield Memorial Lectures VoL II

Fig. 21

Grain boundary reaction in a hypo-eutectoid steel reacted at 750cC for 2.25 h xl00.

Fig. 22

Microstructure in a hypo-eutectoid steel after reaction for 24 h at 700°C xl00.

Fig. 23 Diagram of morphology showing approximate limits of various types of final microstructures observed after isothermal reaction. (1) Area limited by MNP: massive ferrite formed by impingement of growing ferrite crystals previously of Widman statten character or of grain boundary type. (2) Area limited by ABCMNP: Widmanstatten ferrite originated by formation of side plates of ferrite or by nucleation and growth within the grain. (3) area limited by GABSC: grain boundary ferrite, large grains ASTM no. 0-1, small grains ASTM no. 7-8.

The Decomposition

of Austenite by Nucleation and Growth Processes

27

increasing rate of nucleation with decreasing temperature, for which accurate quantitative data are not yet available, is in accord with formal nucleation theory. The growth of these first crystals along the austenite grain boundaries presents a phenomenon worthy of consideration. Figure 24 shows that the growth along the grain boundary of individual ferrite crystals is more rapid than that perpendicular to the grain boundary. It might appear that this effect originates in a more rapid rate of diffusion of carbon along the grain boundary, but this seems to be excluded, for it has been shown that the diffusion coefficient of carbon is not greater at the grain boundary. It must be remembered that the phenomenon of growth may fruitfully be considered as a process in which thin unit layers of the crystal are deposited in succession, that is, as a continuous process of two-dimensional nucleation, a subject upon which Volmer and Stranski have written. Assuming that the same principles apply to the formation of these two-dimensional nuclei as apply to the formation of three-dimensional nuclei, then such two-dimensional nuclei will fonn faster at the grain boundary, that is, growth will be more rapid along the grain boundary.

Fig. 24

Microstructure depicting preferential growth of ferrite along the austenite grain boundaries in a hypo-eutectoid steel. Etched with natal and then with Vilella's reagent x500.

As noted before, the amount of grain boundary pro-eutectoid ferrite varies with the austenite grain size at temperatures below Ael (Fig. 25). At temperatures above Ael the total amount of pro-eutectoid ferrite does not vary with grain size, for it is determined by equilibrium alone. Just below Ae1, both grain boundary ferrite and Widmanstatten ferrite may appear, but the amount of grain boundary ferrite per unit volume is greater in the fine grained steels. At quite low temperatures, where Widrnanstatten ferrite disappears, the thickness of the envelope at each grain boundary is approximately the same in the coarse and fine grained steels, with the result that the total amount of ferrite per unit volume in the fine grained steel is the greater. Thus at high temperature, grain boundary ferrite alone appears; as the temperature of reaction is lower, Widmanstatten ferrite appears, growing as plates from the grain

28

Hatfield Memorial Lectures Vol. II

boundary ferrite; and at still lower temperatures ferrite plates form by nucleation and grow independently within the crystal. Although these Widmanstiitten figures have long won the attention of metallurgists, since the days of Sorby, the mechanism has not yet been wholly clarified. 80~--~--~--~----~--~ a::

2bOI-----+--~~~--t---+------I

~ ~ ~

g401----I---f--~_t__\_t--_I_--_i

z

~201---+--+--__i_+-_+_+_..lIoo~-+--___f

~

Fig. 25

Variation of the percentage of ferrite formed in a hypo-eutectoid steel as a function of reaction temperature.

A comparison of the behaviour of coarse and fine grained steels is instructive. At the same high temperature Widmansrattcn figures will be observed in the coarse grained steel and not in the fine grained. The thickness of the ferrite plates is characteristically large at high temperatures; indeed, in the case cited, the spacing of the plates in the coarse grained steel is greater than the grain diameter itself of the fine grains. It seems that ferrite plates cannot appear as a family of parallel plates in the fine grained steel because the grain diameter is less than the interplate distance; such a steel cannot display the Widmanstatten structure. The spacing of the ferrite plates appears to be quite fundamental. It recalls the spacing between the arms of dendrites, and may well originate in a similar way, modified by the presence of the austenite grain boundary and by the requirement that growth be along the octahedral plane of the austenite. As in the case of dendrites, the interplate distance decreases as the temperature of growth is decreased. In this view, growth proceeds along the grain boundary with a natural tendency to develop crystal faces, and thus edges and corners. Such edges and comers become points at which growth tends to develop preferentially owing to the point effect of diffusion, an effect analogous to the thermal point effect in heat flow. At high temperatures, where growth is slow and the rate of diffusion rapid, there is time for the equalisation of the carbon concentrations between such points and the centre of crystal faces, but at lower temperatures the rate of growth is fast and the rate of diffusion is slow, and equalisation is inadequate, with the result that edges develop and these grow into plates. This will occur periodically during growth, with a resultant spacing determined by how large the crystal may grow before the point effect becomes dominant, and a plate develops. As the temperature decreases, so will the spacing. In the case of dendrites, growth is in the direction of a major crystal axis. These plates grow most rapidly along the octahedral planes of the austenite; owing to the orientation relationship subsisting between the ferrite and the austenite, growth is thus

The Decomposition of Austenite by Nucleation and Growth Processes

29

most rapid in a direction lying in the dodecahedral plane of the ferrite, and least rapid in the direction normal to the most densely packed lattice plane. This appears to be a form of the familiar law of Bravais, originally slated for crystals forming from liquid solutions. The occurrence of ferrite plates independently formed within the grain is obviously a nucleation and growth process of no unexpected characteristics. As in the case of pearlite, the rate of nucleation within the grain is not nil, it is merely very much smaller; given a sufficient time, nuclei will appear within the grain also. It has been shown that the thickening of grain boundary ferrite proceeds parabolically with time. This rate of growth is about the same in both the coarse and the fine grained steels; thus, in the coarse grained steel, more time is available for intragranular nucleation. This reasoning is of the school ofBelaiew, and experiment seems to substantiate it. The curves in Fig. 23 show that the tendency to form Widmanstatten figures decreases at low temperatures. This results from the competitive nature of the two reactions, austenite-to-ferrite and austenite-to-pearlite, as a consideration of the rates of nucleation and growth of the two reactions will show. The rate of nucleation of pro-eutectoid ferrite increases as the temperature decreases, and may actually pass through a maximum in low carbon steels or alloy steels. This nucleation is thus entirely similar in. principle to that of pearlite, but the rate of growth cannot be considered so very simply as in that case. In pro-eutectoid ferrite the concentration of carbon in the unreacted austenite is not unchanging during the reaction, but instead is increasing; the boundary conditions are thus changing during growth. There are two types of initial growth: the growth along the austenite grain boundary, and the growth perpendicular to the grain boundary. The rate of growth along the grain boundary has been measured, as shown in Fig. 26, and appears initially not to vary with time, but after a certain period it decreases. The initial, essentially constant, rate is maintained proportionately for a longer time at lower reaction temperatures. The rate of growth perpendicular to the grain boundary, that is, the rate of thickening of the grain boundary ferrite, is hard to measure, owing to uncertainty in the angle with which the plane of polish intersects the ferrite crystal. But the rate can be calculated, presumably on sound principles. The concentration of carbon in the austenite just ahead

~ ~2.lr--f---t-----.I'--t--i

..----,~-t--_f

Ci

~

OI.rt---4I--+t----+---i

~

~

,

/

1+---+----1 I

~~ Ot.......:....-I-:!:-O-~20=----=:!30 1/ I O=----:-!::IO,..----='20 ~

Fig. 26

TIME. S

TIME,

MIN

Curves depicting the growth of ferrite crystals along the grain boundaries (0.32% carbon, 0.85%

manganese, small grains).

30

Hatfield Memorial Lectures VoL II

of the thickening ferrite plate must follow the type of curve shown in Fig. 27, which is for an infinitely large austenite grain, and will not apply strictly when the concentration at the centre of the grain changes appreciably. At the outset, when grain boundary ferrite first forms, the concentration of carbon in the austenite at the interface immediately approaches the concentration demanded by the extrapolation of the GS curve, and thus the carbon concentration gradient is initially exceedingly great, the rate of diffusion of carbon away from the interface is correspondingly great, and the initial rate of growth is therefore great. This may be seen in Fig. 28, where the composition of the steel is given by the perpendicular X, and the concentration of carbon in austenite at the interface is given by the extrapolation of the GS curve. But as time goes on, the gradient in the austenite diminishes; analysis shows that this provides a parabolic rate of growth, that is, the rate of growth is inversely proportional to the thickness of the ferrite layer. This

Fig. 27 Diagram of the carbon concentration in front of an austenite-ferrite interface: ~ thickness of ferrite plate; Cs1 equilibrium concentration in austenite; Cs2 equilibrium concentration in ferrite; C, interface concentration; C1 initial concentration.

X

Austenite

~800 ex:

:J

<ex:

~700

z:

t=!

b

Ferrite

o Fig. 28

t

cementite

O·S 1·0 I·S (ARBON CONTRATON. WT °/3

Diagram showing the variation of supersaturation with respect to ferrite and cementite as a function of reaction temperature.

The Decomposition of Austenite by Nucleation and Growth Processes 31 applies at all temperatures. Although the calculated rate of growth is difficult to substantiate quantitatively, observation does show that the rate of thickening falls off rapidly with time. It will be observed that the parabolic law of thickening is derived for an infinitely large grain, but when, in a small real grain, the carbon concentration at the centre of the grain decreases, the rate of growth must correspondingly decrease. The parabolic rate constant, however, increases as the temperature falls, for the increasing supersaturation, shown by the extrapolated GS curve, provides a greater concentration gradient, and to a degree sufficient to offset the diminished diffusion coefficient. The rate of growth is a resultant of these two factors, the increasing supersaturation and the decreasing diffusion coefficient. This provides a maximum in the rate of growth at a low temperature when the decrease in the coefficient of diffusion is sufficient to overtake the increase in supersaturation. Such maxima are in fact observed in low carbon hypoeutectoid alloy steels where pro-eutectoid ferrite forms at all temperatures, above and below the pearlite knee. In the discussion of pearlite, the fundamental importance of the ratio of the rate of nucleation to the rate of growth, was commented on, noting that it increased greatly with decreasing temperature, giving more and smaller pearlite nodules; a wholly similar behaviour is observed here, a larger number of smaller ferrite grains occurring at lower temperatures.

The Pearlite Interruption The question of why the formation of pearlite interrupts the growth of grain boundary ferrite at an ever earlier stage as the temperature of reaction is lowered, now needs to be considered. The growth offerrite, and the nucleation and growth of pearlite are competitive reactions; both the rate of growth of ferrite, and the rates of nucleation and growth of pearlite, increase as the temperature is lowered, but the latter increase more rapidly than the rate of growth of ferrite, and the formation of pearlite thus interrupts the growth of ferrite at a proportionately earlier stage. As already stated, the rate of thickening of ferrite can be calculated as a function of the temperature, from the variation of the supersaturation and the variation of the diffusion coefficient with temperature. The supersaturation of the austenite with respect to ferrite is given by the extrapolated GS line in Fig. 28, and with respect to cementite by the extrapolated SE curve. At the interface, where the concentration in the austenite is high, obviously the supersaturation for cementite is high, but at a distance toward the centre of the grain this supersaturation becomes less. Thus the respective supersaturations for cementite and ferrite in the matrix austenite are again given by the extrapolated solubility curves. High supersaturation for cementite implies the imminent precipitation of cementite. It is important to note that this supersaturation is greater than for the nucleation of pearlite in a eutectoid steel, and increases more rapidly as the temperature falls, so that presumably the rate of cementite nucleation does also. Whenever a cementite nucleus forms, pearlite will be nucleated, and the problem of the pearlite interruption is thus the problem of the nucleation of cementite.

32

Hatfield Memorial Lectures Vol. II

The probability of forming a nucleus is proportional to the rate of nucleation and to the time available for nucleation. When grain boundary ferrite first begins to form, the ferrite-austenite interface is moving rapidly, that is, there is not much time for nucleation. As the rate of growth falls off with time, the probability becomes greater, and ultimately a cementite nucleus forms. Once pearlite is nucleated by this cementite, it grows at a characteristically high rate as compared with the rate of thickening of the ferrite at the grain boundary; starting in the austenite contiguous to the ferrite where the carbon concentration is high, the cementite plates are thick (Fig. 29) and as growth continues into lower carbon areas, the cementite plates become thinner.

Fig. 29

Microstructure illustrating the thickening of cementite plates at a pearlitic-ferrite interface. Etched with picral x2000.

It can be shown from these considerations that the thickness of the ferrite at which the nucleation of pearlite interrupts further ferrite growth, decreases as the temperature decreases. This interruption of the growth is obviously the factor which prevents the formation of Widman statten plates at low temperatures. The interruption of the thickening of the ferrite at an increasingly earlier stage on decreasing temperature probably occurs because the supersaturation at the interface at low temperature becomes so large that the cementite nucleus may form even though the thin grain boundary ferrite grows at a rate which is characteristically high. It will thus be obvious that the value of the supersaturation at the interface, and the velocity of the thickening of the ferrite grain boundary layer govern the point of interruption by pearlite nucleation, affording a critical thickness of the ferrite layer at each reaction temperature; this critical thickness will be the same in steels of different grain sizes, providing more ferrite per unit volume in the fine grained steel than in the coarse grained steel. In steels of different carbon contents, and at different temperatures, the carbon composition of pearlite thus changes, and the rate of growth changes. Measurements of the rate of growth of pearlite in steels with 0.32 and 0.62% of carbon, respectively, show that the rate of growth is the greater the lower the carbon content; this is entirely in accord with Brandt's calculations.

The Decomposition

of Austenite by Nucleation and Growth Processes

33

The isothermal rate curve for this process is in two parts, representing successive reactions. In the early stages, where only grain boundary ferrite fonns, such curves can be calculated, employing a rate of nucleation and a rate of growth decreasing parabolically with time, but when Widmanstatren ferrite forms, the geometry becomes very complex and the reaction rate equation also. The second stage, when pearlite forms in a matrix interlaced with ferrite plates, is similarly complicated geometrically, for this ferrite hinders the growth of pearlite, providing steric hindrances resulting in a much slower rate. This steric factor thus affects the pearlite rate to the degree that the pro-eutectoid ferrite forms plates. It is implicit in the foregoing that the ferrite and the pearlite reactions must overlap to some extent, that is to say, ferrite continues to form at one point for a while after pearlite has already started to form at another point. The nucleation of pearlite following the formation of ferrite is a nucleation phenomenon, and thus governed by probability. If ferrite thickened with utmost precision, forming no protuberances, providing a ferriteaustenite interface wholly plane, the initiation of the pearlite reaction would be governed by probability alone, providing a certain small degree of overlap, but this is augmented by the variations in concentrations in the austenite resulting from irregular ferrite growth. The degree to which overlapping occurs has not been studied systematically. Recent work has shown that hypereutectoid alloys behave similarly in many respects. Similar rules appear to apply to the nucleation and growth of the grain boundary phase, and of the Widmansratten figure, but cementite plates grow from grain boundary cementite only infrequently, and Widmanstatten precipitation by separate nucleation and growth within the grain is far more pronounced, forming a sharply defined Widmanstatten figure of thin long plates.

PRO-EUTECTOID

REACTIONS IN ALLOY STEELS

Pro-eutectoid reactions in alloy steels, i.e. the formation of pro-eutectoid ferrite and proeutectoid cementite, show many similarities to the corresponding reactions in carbon steels. The reaction is again a nucleation and growth process, the pro-eutectoid constituents forming preferentially at the austenite grain boundaries, the ratio of the rate of nucleation and growth increasing as the temperature of reaction is decreed and Widmanstatten figures again appearing. Although these generalisations are quite justified, there have been no careful studies on morphology and none on rates of nucleation and growth; only scattered photomicrographs and measurements of the overall reaction rate are available. The addition of alloy may well modify both morphology and rates of nucleation and growth, for obviously the diffusion phenomena are more complex, the possibility of concentration of the alloying element at the austenite grain boundary appears here again, and the interface energies must be altered; these factors must affect both morphology and rates. The field is open for definitive study. It was observed in considering alloy eutectoid steels that in some steels pearlite forms both above and below the pearlite knee; such alloy steels differ in th~s way from carbon

34

Hatfield Memorial Lectures VoL II

steels and from nickel and manganese alloy steels in which pearlite forms only at temperatures above the knee. This behaviour of pearlite in alloy eutectoid steels has its counterpart in pro-eutectoid ferrite in low carbon alloy steels, ferrite forming at temperatures both above and below the temperature of the knee, and when the carbon is sufficiently low, exhibiting a knee of its own. This knee, as in the case of pearlite, must indicate a maximum in the rate of nucleation or in the rate of growth, or both, and may find its explanation in the competitive processes of free energy change and diffusion rates, a possibility anticipated in formal nucleation theory. Measurements of these rates are necessary before the problem can be formulated. The phenomenon of a ferrite knee in the isothermal diagram is exhibited by alloy steels in which the alloying element is carbide forming, for example, chromium, and these steels constitute a class in this respect. Other alloy steels, as exemplified by nickel steels, form a second. class, in which the alloying element does not change the form of the isothermal diagram, but merely displaces it to longer times. In such steels pearlite forms only at temperatures above the knee. The correlation of this behaviour with the carbide forming propensities of the alloying element is good, but the basic reason for the distinction of the two classes of steels has escaped formulation. It would be comforting if full and satisfactory knowledge of the equilibrium constitution of alloy steels were available, for this would provide some foresight into the phase constitution to be expected, and would illumine departures from equilibrium as determined by kinetics, but information on the constitution of even the ternary systems of iron is very limited, and it is usually impossible to define the final equilibrium state. Although but little work has been done on the composition of pro-eutectoid ferrite in alloy steels, beautifully executed work by Bowman on hypo-eutectoid molybdenum steels has shown that pro-eutectoid ferrite in these steels contains exactly the same concentration of molybdenum as the austenite from which it forms, i.e. molybdenum does not diffuse during the formation of ferrite. In view of this surprising though well attested result, it is difficult to understand why molybdenum, like all other alloying elements, retards the formation of ferrite. Blanchard and others believed that the slight effect of molybdenum in decreasing the rate of diffusion of carbon might be the cause, though this hardly seems adequate, or that molybdenum might change the limiting concentrations in such a way as to decrease the diffusion gradient, or, finally, that molybdenum might actually retard the fundamental ,,{-u transformation, i.e. that the interface reaction itself, apart from diffusion, might in this case be rate determining. The subject is difficult to pursue at this time, for apart from these molybdenum steels, there is little useful information; much more work on other systems, with the same end in view, is greatly needed. Observations on the effect of the rejection of pro-eutectoid constituents upon the subsequent rate of formation of pearlite, show that molybdenum not only retards the formation of ferrite, but also so alters the kinetics of the system that the pearlite reaction ineterrupts the growth of ferrite at a later time, that is, molybdenum increases the amount of pro-eutectoid ferrite. The same effect has been observed by Lyman and Troiano over

The Decomposition

of Austenite by Nucleation and Growth Processes

35

the entire carbon range in chromium steels. Pro-eutectoid cementite in molybdenum steels has the opposite effect. Inthis case cementite actively nucleates pearlite, for which it is a natural inoculant, thus decreasing the time for initiation of the formation of pearlite, and grain boundaries do not play their usual role in controlling the number of nuclei formed. In all such successive reactions, kinetics are still determined by nucleation and growth, but when the pro-eutectoid constituent forms plates, the subsequent pearlite grows with continued hindrance, the steric factor again being prominent. A reaction rate equation would have to be written with this steric factor properly included. Alloying elements may thus produce a ferrite knee in the isothermal diagram, which anticipates a pearlite knee occurring at the same temperature, with both ferrite and pearlite forming at temperatures both above and below the knee, as with the steel containing 1% of chromium and 0.4% of molybdenum studied by Klier and Lyman. Whereas in carbon eutectoid steels pearlite forms only at temperatures down to the knee of the isothermal diagram, and bainite below, in the class of steels exemplified by chromium steels, pro-eutectoid ferrite forms below the knee of the pearlite curve, down to the region where the X constituent and bainite form, providing an intermediate temperature at which the rate of reaction is very slow (Fig. 5). It is proper now to consider the complex case of bainite.

BAINITE Below the point of the knee of the isothermal diagram in eutectoid carbon steels, and in alloy steels in the region of the 'bay'· in the. isothermal diagram, new products of the decomposition of austenite appear, differing fundamentally in nature from pearlite. These structures, which will be referred to by the generic term 'bainite', and which are formed at intermediate or low temperatures, are by no means new, for they were observed early in this century. The structures are simplest as they occur in eutectoid carbon steels. As formed at high temperatures (Fig. 30) the structures are coarser than low temperature pearlite, that which used to be called nodular troostite. The progress of the transformation appears to be the thrusting of thin fingers of ferrite into the austenite matrix, followed by the precipitation of carbide, providing a feathery structure. As Hultgren pictures it (Fig. 31) ferrite forms first, momentarily enriching the austenite locally in carbon to a point where carbide, precipitates in particles, with ferrite growing between. The ferrite plates grow large because of the relatively high rate of diffusion of carbon and the relatively low value of the rate of nucleation of the cementite. As the temperature of isothermal decomposition is lowered, the structure changes gradually until reaction just above the martensite point produces a dark etching acicular structure difficult to distinguish from tempered martensite (Fig. 32). These structures have proved difficult to study. Observations by G. V. Smith indicate that both the number and the size of bainite patches in the resultant structure increase as

36

Hatfield Memorial Lectures Vol. II

Fig. 30

'Upper bainite' for eutectoid steel partially reacted at 540°C (ViIella) x1500.

Jl Fig.31

Fig. 32

Hultgren's model for the formation of bainite.

'Lower bainite' in eutectoid steel formed upon reaction at temperatures near the martensite range (Vilella) x2500.

the reaction proceeds, clearly suggesting a nucleation and growth process, and suggesting an attack similar to that employed on pearlite, the measurement of the rates of nucleation and growth. Nothing has been done in this direction, as this part of the study of the fundamentals of the decomposition of austenite, like the rest, received but little attention during the war years.

The Decomposition of Austenite by Nucleation and Growth Processes 37 If it is assumed, on the slight evidence that exists, that bainite does form by nucleation and growth, rather than, as in the case of martensite, by shear, the alternative process, then it becomes evident immediately that the morphology of the product imposes special steric factors: the advancing front (Fig. 30) is not plane and the reaction is not consummated at that interface, and the structure frequently shows interlacing, so that the overall rate of decomposition, though fundamentally dependent upon rates of nucleation and growth, can be calculated from these only by a very complex and possibly fruitlessly complex formulation. The problem is therefore not unlike that in the formation of proeutectoid ferrite, where the geometrical complexity in the disposition of the reaction product makes analysis of reaction kinetics difficult. In the case of high temperature bainite (and in many alloy bainites) the structures formed are exceedingly similar to those in pro-eutectoid ferrite. Bainite forming along austenite grain boundaries throws plates of ferrite toward the centre of the grain, and these plates appear to be mere outgrowths from the grain boundary ferrite. This behaviour is never seen in martensite structures. Nucleation

and Kinetics

The best evidence on the question of nucleation is crystallographic. The relationship between the orientation of the ferrite in bainite and the austenite from which it forms in eutectoid carbon steels is identical to that observed when ferrite is known to nucleate directly from austenite, as in the formation of pro-eutectoid ferrite in hypo-eutectoid steels, as G. V. Smith showed. The conclusion seems inescapable that the active nucleus in the formation of bainite is ferrite; in this way the bainite reaction is to be distinguished fundamentally from the pearlite reaction. These orientation relationships between ferrite and the parent austenite are observed throughout the temperature range in which bainite forms; the orientation .relationship subsisting between two lattices in space, and the plane in austenite to which the reaction product in the form of plates is observed to lie parallel, are to be clearly differentiated. This plane, the composition or conjugate plane, varies throughout the temperature range even though the orientation relationship remains constant. The phenomena of growth are in truth often complex, and in this case wholly unrationalised. Although there is no supporting evidence, it has been assumed universally that the ferrite nucleus in the formation of bainite is ferrite supersaturated with carbon, presumably because the ferrite which forms precedes in point of time the formation of cementite. This supersaturated ferrite is conceived as ephemeral, precipitating carbide very quickly; the argument is awkward, for it is surely not proof, and a direct experimental attack, difficult to conceive, would be very welcome. In eutectoid carbon steels the pearlite reaction, proceeding more and more rapidly as the temperature decreases, is replaced at the knee of the isothermal diagram by the formation of bainite, which forms more and more slowly as the temperature decreases. The knee of the isothermal diagram is thus a point where one reaction is replaced by a fundamentally different one. There is no sensible break in the isothermal diagram at this

38

Hatfield Memorial Lectures Vol. II

point in eutectoid carbon steels, as one might expect. Presumably, extremely precise measurements would show a discontinuity; some overlap of the two reactions is always observed with samples reacted just at the knee showing both pearlite and bainite, but how much of this is to be ascribed to austenite heterogeneity is unknown. However, in some alloy steels, those in which the alloy is strong in carbide fonning proclivity, the separation of the two processes is very evident. The isothermal diagram for many alloy steels shows, as noted above, a full knee for the pearlite reaction, and a second, lower knee for the bainite reaction, separated by a bay of extremely slow reaction. Although other proposals have been made, it seems most likely that the cessation of the pearlite reaction is brought about by the intrusion of the faster bainite reaction. This argument is a little, though not wholly, truistic. To develop it would require better knowledge of the true nature of the nucleus, whether supersaturated ferrite or not, and a sustaining knowledge of the factors which determine the rates of nucleation and growth for both 'pearlite and bainite, and how these might be modified by alloying elements. Here, again, existing knowledge is nearly uselessly rudimentary. Zener has proposed that the upper temperature limit of bainite, the temperature below which bainite can form, is determined by the free energy change accompanying the formation of a ferritic phase of the same composition as the original austenite, though there have been some objections to this. The, formation of martensite he considers is quite similar, except that in this case additional free energy is required, and thus a lower temperature, in order to overcome the accompanying strain energy.

Bainite in Alloy Steels The formation of bainite in alloy steels presents a diversity of phenomena that is perplexing. This field, of the intermediate temperature transformation products in alloy steels, is very important practically, for these products occur widely in heat treated steels. The lack of knowledge can be ascribed to several circumstances: the fundamental kinetics of nucleation and growth are but embryonic; experimentally the microscope has been relied upon too much (a greater diversity in the experimental attack on the problem is needed); and the work done has notbeen systematic enough (too many miscellaneous unrelated steels have been studied, though this is not true of the work on chromium steels and molybdenum steels). Little can be said on this subject beyond relating several of the simple features which are well attested. The low temperature bainite knee in many alloy steels is attended by the initial rejection of a ferrite like constituent, usually acicular in character, which has been called the 'X constituent' (Fig. 33) and which is sometimes also referred to as bainite, probably inadvisedly. The formation of this constituent is well illustrated in silicon steels. It has been believed that this structure when first formed is ferrite supersaturated with carbon, but its nature is not well understood, and the relationship of this structure to bainite formed at similarly low temperatures in other alloy steels also is not understood. In 3% chromium, 0.38% carbon steels, Lyman and Troiano observed a bainite to form at 490°C

The Decomposition of Austenite by Nucleation and Growth Processes

39

at nearly zero time, but that only 10% of the austenite reacted to this constituent, the reaction then nearly stopping, succeeded in time by the formation of pearlite. At 450°C this reaction again starts at nearly zero time, but proceeds to 80% of completion, whence the remainder of austenite forms pearlite. This temperature dependence of the reaction to bainite suggests the similar behaviour of martensite, but the analogy should not be forced, for the bainite reaction proceeds isothermally with time, whereas martensite does not form isothermally.

Fig. 33

'X constituent' formed in steel with 0.16% carbon, 3.36% nickel, 0.52% molybdenum, reacted in upper region of bainite range (ViIella) x2000.

Hultgren's fine paper illustrates many of the phenomena and therefore the complexities of the alloy bainites. Arguing that high temperature products in the pearlitic region are close to equilibrium, with both carbon and alloying elements migrating, and that in martensite neither carbon nor alloying element migrates, he believes that at intermediate temperatures products should form from which carbon has migrated but not the alloying elements. The high temperature products he designates as orthoconstituents and the low temperature products as paraconstituents. He observes, indeed, that the carbides formed in the bainitic range are not at equilibrium as shown by residue analysis, for example, a 3% manganese, 0.52% carbon steel provides a bainitic carbide of the same composition in manganese as the austenite from which it formed. Assuming from this that bainitic ferrite is similarly far from equilibrium, that is, much too high in manganese, he conceives bainite as constituted of paraferrite and paracementite. Although the evidence is scanty, the suggestion is important, for the phenomena of variations in composition of both the reactant austenite and the resultants ferrite and carbide are important and frequently inscrutable. The difficulty in this matter is chiefly the lack of weighty evidence, and the difficulty of getting it, and as much the lack of knowledge of true equilibrium in these polycomponent systems. Griffiths, Allen and Pfeil, in two admirable papers, published by this Institute (The Iron and Steel Institute), under the sponsorship of the Alloy Steels Research Committee, of which Dr Hatfield was

40

Hatfield Memorial Lectures VoL II

Chairman, studied many of these questions, observing that austenite grain size has but little, if any, effect upon the rate of formation of bainite. The complexities in behaviour which are encountered in the bainitic reaction are somewhat awesome. Lyman and Troiano, studying a 3% chromium, 1% carbon steel reacted in the range 350-450°C, observed that the reaction proceeds by separate carbide, ferrite and pearlite reactions, and that cementite plates may form after partial decomposition to bainite, while Hultgren, working with a 5% manganese, 0.63% carbon steel, found pro-eutectoid ferrite to form above the pearlite knee, and pro-eutectoid cementite below. Morphology also is complex: Hultgren observed in a steel with 2.4% of manganese and 1.02% of carbon that on reaction at 500°C cementite plates are formed, and then bainite, with the carbide from the bainite attaching itself to the primary carbide, producing a degenerate bainite. Blanchard, Parke and Herzig, observing the behaviour of a 0.77% molybdenum steel reacting at 538°C, noted that the primary reaction is the formation of ferrite in plate form lying parallel to the octahedral Widmanstatten plane, and that this reaction is followed by the formation of cementite in plate form lying parallel to its Widmanstatten plane. Lyman and Troiano observed that the sequence of reactions at a low temperature in a 3% chromium, 0.69% carbon steel, is the formation of ferrite followed by a rejection of cementite, and this in tum is followed by the formation of pearlite at 450°C; and bainite, followed by cementite, and this in tum by pearlite at 400°C (Fig. 34).

Fig.34

Isothermal transformation diagram for a 0.69% carbon, 3% chromium steel (Lyman and Troiano).

The complexities in bainite, and the lack of basic, guiding ideas, but illustrate in a magnified way the needs in this whole field. The phenomena in this field for which only a beginning in understanding has been made, are of basic importance to the iron and steel industry, and although the discussion bears the guise of science, the matters discussed are in the ultimate view practical in nature, for they underlie practice in the handling of steels. Experiment in this field should be more systematic in nature. It would be better,

The Decomposition

of Austenite by Nucleation and Growth Processes

41

the author believes, especially in forwarding a knowledge of bainite, to take one ternary alloy steel for study, and to determine the equilibrium relationship with great care; to study the morphology of the products, and particularly to attempt to correlate the morphology of pro-eutectoid ferrite with that of bainite; and to study the composition of the separate constituents, so that the enrichment of one phase at the expense of another may more nearly be defined. This latter point is important: a general method for the study of the constitution of micro constituents is still lacking; the chemical analysis of carbides by residue, a partial solution, is unsatisfactory, and has been very little forwarded in recent years. With proper attention to electrochemistry and with less conservatism in the selection of reagents for chemical attack, marked advances are surely possible; but some more general method also is needed to study the composition of constituents. If the use of radio isotopes can be developed for determining the partition of an element between two constituents in structures that require high magnification for their resolution, then marked progress could be made. Morphology, in part determined by kinetics and in part determining kinetics, ought to be correlated with the great body of chemical information on the nature of crystal growth. The opportunities of cross-fertilisation in this field, as in kinetics, are apparently great. Nucleation theory is patently a branch of the general field of kinetics, and ought to be developed as such a branch; it needs to be extended toward a greater degree of reality in application to solid-solid reactions, and what is needed in this is the measurement of energies at phase interfaces, and, just as important, a general and direct attack upon the thermodynamics of these reactions, particularly a method of determining the free energy changes that accompany the many complicated reactions which are encountered. The metallurgist brings very real talents to bear on these problems, and with further attention to fundamentals he should engage himself in the quest of converting this ancient art of heat treatment to a new science.

BIBLIOGRAPHY ]. O. ARNOLDand A. MCWILLIAM:'The Thermal Transformations of Carbon Steels',]. Iron Steel Inst., 1905,2,27. H. MIERS and F. ISAAC:'The Spontaneous Crystallization of Binary Mixtures-Experiments on Salol and Betol.' Proc. R. Soc. A, 1907,79, 322. N. T. BELAIEW:'On the Genesis of Widman statten Figures in Meteorites and in Iron-Nickel and Iron-Carbon Alloys', Mineralogical Magazine, 1924, ,20, 173. E. S. DAVENPORTand E. C. BAIN: 'The Transformation of Austenite at Constant Subcritical Temperature.' Trans. Am. Inst. Mining Metall. Eng. 1930,90, 117. G. KURDJUMOWand G. SACHS:'The Mechanism of the Hardening of Steel.' z. Phys., 1930, 64, 325. F. WEVERand others: A series of papers pubished in Mitteilungen aus dem Kaiser- Wilhelm- Institut fur Eisenforschung zu Dusseldoif, 1930.

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Hatfield Memorial Lectures Vol. II

H. C. H. CARPENTER andJ. M. ROBERTSON: 'The Formation

of Ferrite from Austenite.'J.

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of Austenite

M. VOLMER: 'Kinetics of the Formation C. WELLS and R. F. MEHL: 'Rate Manganese,

on Cooling;

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]. R. BLANCHARD, R. M. PARKE and A. J. HERZIG: 'The Isothermal,

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M. A. GROSSMANN and R. L. STEPHENSON: 'The Effect of Grain Size on Hardenability',

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Trans. Am. Soc. Met., 1941,29,317. Am. Soc. Met., 1941, 29, 1. R. F. MEHL: 'The Structure

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Trans. Am. Soc. Met., 1941,29,813. F. C. HULL and R. F. MEHL: 'The Structure

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F. C. HULL, R. A. COLTON and R. F. MEHL: 'Rate Pearlite',

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Trans. Am. Inst. Metal Eng. 1942, 150, 185.

G. E. PELLISSIER and others: 'The Interlamellar 29,1049. G. V. SMITH and R. F. MEHL: 'Lattice Martensite'

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Trans. Am. Inst. Mining Metall Eng. 1942, 150, 211.

Trans. Anl. Soc. Met. 1942, to Pearlite,

Bainite

and

The Decomposition of Austenite by Nucleation and Growth Processes

43

'Atlas of Isothermal Transformation Diagrams', The United States Steel Corporation, New York, 1943. ]. R. BLANCHARD,R. M. PARKEand A. ]. HERZIG: 'The Effect of Molybdenum on the Isothermal, Sub-critical Transformation of Austenite in Eutectoid and Hypereutectoid Steels', Trans. Am. Soc. Met., 1943,31,849. F. E. BOWMAN,R. M. PARKEand A.]. HERZIG:'The Alpha Iron Lattice Parameter as Affected by Molybdenum, and an Introduction to the Problem of the Partition of Molybdenum in Steel', Trans. Am. Soc. Met., 1943,31,487. J. L. HAM,R. M. PARKEand A. J. HERZIG:'The Effect of Molybdenum on the Rate of Diffusion of Carbon in Austenite', Trans. Am. Soc. Met., 1943,31,877. F. E. BOWMAN and R. M. PARKE: 'The Partition of Molybdenum in Iron-CarbonMolybdenum Alloys at 1300 Degrees Farhrenheit, and the Nature of the Carbides Formed', Trans. Am. Soc. Met., 1944,33,481. E. P. KLIERand T. LYMAN:'The Bainite Reaction in Hypoeutectoid Steels', Trans. Am. Inst. Mining Metall. Eng., 1944, 58, 394. W. A. ANDERSONand R. F. MEHL:'Recrystallisation of Aluminium in Terms of the Rate of Nucleation and the Rate of Growth', Trans. Am. Inst. Mining Metall. Eng., 1945. 161, 140. F. E. BOWMAN:'Partition of Molybdenum in Hypo-eutectoid Iron-Carbon-Molybdenum Alloys', Trans. Am. Soc. Met., 1946,36,61. F. E. BOWMAN:'The Partition of Molybdenum in Steel and its Relation to Hardenability', Trans. Am. Soc. Met., 1945, 35, 112. W. H. BRANDT: 'Solution of the Diffusion Equation Applicable to the Edgewise Growth of Pearlite',]. Appl. Phys., 1945,16, 139. ]. L. HAM: 'The Rate of Diffusion of Molybdenum in Austenite and in Ferrite;, Trans. Am. Soc. Met., 1945,35,331. E. P. KLIER:'Transformation of Austenite in a Steel Containing Three Percent Chromium and One Percent Carbon;, Trans. Am. Inst. Mining Me tall. Eng., 1945,162,186. T. LYMANand A. R. TROIANO:'Isothermal Transformation of Austenite in One Percent Carbon, High Chromium Steels', Trans. Am. Inst. Mining Metall. Eng., 1945. 162, 196. W. H. BRANDT:'Some Factors Affecting Edgewise Growth of Pearlite', Trans. Am. Inst. Mining Metall. Eng., 1946,167,405. T. LYMANand A. R. TROIANO:'Influence of Carbon Content Upon the Transformations in 3 Percent Chromium Steels', Trans. Am. Soc. Met., 1946,37,402. E. SCHElL:'The Eutectic Crystallization', Metallforschung, 1946, 1, (1 and 2), 1. E. SCHElL:'The Calculation of the Velocity of Eutectic Crystallization with Pearlite as an Example', Metallforshung, 1946, 1, 123. C. ZENER: 'Kinetics of the Decomposition of Austenite', Trans. Am. Inst. Mining Metall. Eng., 1946, 167,550. M. F. HAWKESand R. F. MEHL: The Effect of Cobalt on the Rate of Nucleation and the Rate of Growth of Pearlite. Technical Publication No. 2211, American Institute of Mining and Metallurgical Engineers, 1947. A. HULTGREN:'Isothermal Transformation of Austenite', Trans. Am. Soc. Met., 1947,39,915. A. R. TROIANOand]. E. DEMoss: 'Transformations in Krupp-Type Carburizing Steels', Trans. Am. Soc. Met., 1947,39,788.

44

Hatfield Memorial Lectures Vol. II

A. DUHEand R. F. MEHL: Unpublished research on the Decomposition of Austenite in HypoEutectoid Steels. R. W. PARCELand R. F. MEHL:Unpublished research on the Influence of Molybdenum on the Rate of Nucleation and the Rate of Growth of Pearlite. C. WELLSand W. BATZ:Unpublished research on the Diffusion of Carbon in Austenite.

THE

EIGHTH

HATFIELD

MEMORIAL

LECTURE

Trends in Metallurgical Research in the United States Edgar C. Bain At the time the lecture was given Dr Bain was Vice-President, Research and Technology, of the United States Steel Corporation and member, National Academy of Sciences. The lecture was presented in the Firth Hall of Sheffield University on Monday 17th October 1955.

For the honour of being invited to present the Eighth Hatfield Memorial Lecture I am deeply appreciative. I am grateful, too, for the distinction of having my name added to the list of illustrious former Lecturers who are thus associated with the memory of Dr W. H. Hatfield. I recall well his visit to the American Society for Metals in 1928 as the Edward deMille Campbell Lecturer; he came to us - characteristically, I suppose with a manuscript for a book of 150 pages, using only small excerpts therefrom for the Campbell Lecture. The book, later published, bears the title The Application of Science to the Steel Industry. A fairly significant lecture for this occasion might be written, if one had the courage, just on the provocative questions which Dr Hatfield raised or implied in this 1928 book, particularly if we added the current answers - if any which we might make today. I remember also the enthusiastic, congratulatory slap on the back he gave me in his laboratory, in 1932, when, by way of a demonstration, I bent a small round bar of our 'austempered' steel (British Pat. No. 424,124) around approximately its own diameter, though it was hard enough to scratch glass. Dr Hatfields's 1928 visit to America came in the midst of a period during which metallurgists in general there were beginning to feel some confidence in their own competence in metallurgical research. Previously, many of the more refined and brilliant studies had been made in Europe, and there had been a feeling that from Europe must come the fruits of any very difficult researches, such as constitution diagrams and the like. It was a period, too, in which light was being cast rather easily upon many problems, and we felt, perhaps, that we were reaching a kind of steady state in which what we called the 'science of metals' would grow at a comfortable, constant rate through our now effective researches. What with excellent microscopy, dilatometric equipment, X-ray diffraction apparatus, etc., in competent hands, I should conjecture that what we were really doing was, to some extent, picking all the easy tasks and making a rather good job of them. Perhaps we shall never see an era again of such apparent lucidity of theory, and of such so-called 'down-to-earth' experimentation. To be sure, we did inveigh against designing with built-in notches, but no one then lost much sleep over high velocity, brittle fracture or transition temperature. We were, we felt, just about to round out the last item in the full

45

46

Hatfield Memorial Lectures VoL II understanding of martensite. Nucleation and growth had been a very appropriate descriptive phrase, and given a simple quantitative significance by Tammann. The subject was yet to receive its detailed extensions under R. F. Mehl. And with something approaching nostalgia, I remind you how crystals of metals were then fully represented by cork balls on rods; and in those 1928 models there were no vacancies provided, nor any edge dislocations. If I may be forgiven a reminiscence, it was during 1928, twenty-seven years ago, that I spent my spare time wondering how one might somehow learn something quantitative about the time consumed in the direct transformation of austenite in steeL I had already made some semi-quantitative observations upon the isothermal transformations of retained austenite, retained, persistently, that is, in untempered, quenched tool steels of fairly high alloy content. I decided early that I should have to study the rates isothermally if I were ever to unscramble the data. It was pleasant and fortunate for me, personally, that Mr E. S. Davenport in 1929 decided to cast his lot with me in this venture, and the work prospered thereby.t-f The ensuing study was rewarding in the kind of pleasant satisfactionwhich research can give. Our vast regret is that our preoccupation with the distinct types of microstructure formed, and their growth patterns, especiallyin the temperature regions of high time interval accuracy, led us so long to neglect, and thus misinterpret, the rates and mechanisms observed in the lower temperature zone of the isothermal transformation diagrams. In written discussion at the time, Robertson.' pointed out that martensite forms, not isothermally, but athermally with falling temperature. This was an amazing concept; it required time for confirmation and assimilation, but at length we all recognised its verity and significance. The most unfortunate part of any error which is published is that it manages to be copied and re-published so often before the correct view is widely promulgated. We sought finally to correct the older and unacceptable story in the second edition" of the Atlas of Isothermal Transformation Diagrams, which, with the supplements, now embraces 474 compositions. In the Atlas, the lower part of the diagram is omitted as not being a range for isothermal transformation. The British Atlas is, I believe, modelled on the United States Steel hook, though your volume escaped the error most thoroughly. And now, ironically enough, Morris Cohen> at Massachusetts Institute of Technology finds that there are isothermally formed martensites in certain compositions. But we reminisce too long.

We are today to discuss trends in metallurgical research in the United States. It will, I am sure, cover only a few of many trends, for research, at home as elsewhere, is like the cavalier who mounted his steed and galloped off in all directions; nor do I think that this is necessarily all bad. In the case of some of the directions taken by metallurgical research, I am far from qualified to discern, much less describe, a trend. At the outset I can say that in our thinking, if not in our budgets, the old scientific distinction between ferrous and non-ferrous metallurgy is not held to be very significant or important. Rather, iron is counted as one of the more interesting metals, with somewhat more complicated behaviours, and hence we not infrequently find it

Trends in Metallurgical Research in the United States

47

convenient first to study other metals, the better thus to approach the peculiar problems of iron and steel. One trend in our fundamental researches of late will, I predict, prove of great advantage in the growth of our 'science of metals.' I refer to the recent tendency to combine critical experimental programmes with the theoretical creativeness which has flourished so outstandingly on both sides of the Atlantic during the last few years. In other words, there appears to be a new desire to return to ingeniously controlled experimentation designed to confirm or deny a theory, or just to add significant understanding of some metallic behaviour. This trend is illustrated well in the collected papers of the nine pre-Congress Seminars" held over the weekends preceding the regular technical sessions of the American Society for Metals. The papers usually cover some timely, lively, highly scientific and engrossing subj ect and they are thoroughly discussed. These seminars are an inspiration for the younger technical people and, for the older, a source of amazement. I should be inclined to give a very high place to these meetings in the furthering of fundamental metallurgy. In the matter of epistemology, there seems to be a strong, though gradual, recognition in the United States of the consequences of what may prove to be an unfortunate categorisation. This perplexity has to do with metallurgy itself and its metes and bounds. One needs from time to time to reflect upon the third syllable, i.e. the '-urg-' of metallurgy (which we, in our American pronunciation, accent rather more strongly); it is, like 'erg,' from ergos, that is, a 'worker.' A metallurgist was, then, philologically, a worker in metals. By extension, is a research metallurgist one who works at research in metals or one who carries on research in metal working? More seriously, the difficulty which arises may have its origin, in part, with our present concept of metallurgy as a science of metals, instead of the working of metals as a skill. Traditionally and conventionally, the branches of science are quite properly named and defined to represent domains of scientific knowledge and, thus categorised, a satisfactory degree of homology or coherency exists within each. But the name 'metals' (including alloys) applies to a certain class of substances, not a domain of knowledge. The scientific basis for understanding metals is derived from many of the branches of science, and its domains cut across nearly all. Fortunately, one need be highly trained in only two or three of the various sciences to study and understand much of the behaviour of metals and a very great deal of their technology. Indeed, that is what is largely, and so wisely, taught in courses in metallurgy. But the philosophy of categories is at times stretched far to embrace 'the science of metals.' The curious circumstance emerges that the particular scientist, who, from one entirely sound point of view, knows most of what a bit of metal is really composed, is probably utterly unable to aid in producing, for example, a better heat treated forging. I mention this in particular because we are facing a grave difficulty in securing sufficient metallurgists either for operations or for research. Actually, it would appear that we should rather say that we are experiencing a dearth of chemists, physical chemists, chemical engineers and physicists, schooled in the special disciplines of the solid metallic

48

Hatfield Memorial Lectures VoL II

state, who wish to make careers in metallurgy. I venture to hope that students are not repelled from metallurgy by the third syllable, '-urg-' which may perhaps suggest ergon, rather than ergos, and means 'work.' Turning now from my own musings on the anatomy of metallurgy as an exceedingly broad segment of science, or a confluence of sciences, I should like to refer specifically to some of the metallurgical researches which are under way in America, and some of the generalisations which are taking form. They are not necessarily the most important, the most interesting or even the most recent. They are merely a selection.

X-RAY METALLOGRAPHY AND RELATED TECHNIQUES Great impetus has been given to several significant applications of X-ray diffraction by the use of electronic counters replacing photographic film. As shown in the comparison of intensity measurements over a broad line, in Fig. 1, the counter can significantly improve the techniques. Enhancing the accuracy of intensity and angle measurements as indicated, the Geiger counter, and for wider linear range of intensity response, the proportional or scintillation counters, have extended X-ray applications widely. Furthermore, the automatic recording of angle and intensity, by conserving time, has greatly expedited studies of the constitution of alloys. The work of Goldschmidt? at the BSA Laboratories in Sheffield is regarded by American investigators as setting the highest standards in its field.

f

Photographic

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DI FFRACTION

ANGLE

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Comparison of intensity distributions about the a1 peak of the (211) reflection from ferrite in quenched and tempered steel as measured by Geiger counter and photographic techniques. AISI 4342 steel oil quenched from 1550°F, tempered 0.5 h at 40QoF, air cooled. Harness 53.5 Cr radiation 40 kV, 10 rrtA.

Fig. 1

u;

tc;

Trends in Metallurgical Research in the United States 49 Attempts have long been made to use X-ray diffraction to measure elastic strainf and therefrom to evaluate stress (particularly residual) near surfaces. The elastic anisotropy of iron (and certain other metals) greatly complicates the problem. Dr John Norton? says the real question is one of whether the particular grains which the X-ray beam selects from the poly-crystalline specimen correctly represent the true average macroscopic stress, or whether, because of the high level of heterogeneous microstresses, something different from the true average macrostresses is indicated. Probably, nearby grains are in neither an 'iso-strain' nor an 'iso-stress' state. The precision method of Norton!" employs the apparatus shown in Fig. 2. It involves getting the difference in spacing of chosen planes, as between those very nearly parallel to the surface and a set of corresponding planes at an angle of about 45° thereto, both measurements in the same plane with reference to the principal stress direction in question. A calibration table, made up from data on originally stress free material under known macrostress externally applied, provides the values of stress from the X-ray crystal strain data. Similar surface and 45° measurements are often made at right angles to the first, thereby very largely defining, for example in plate, the state of stress at the surface. As may be seen in the exploration of longitudinal stress in a notched plate specimen (Fig. 3), consistent results may thus be obtained with high precision; but such achievement involves all devices for pushing the method to its very limit. Again, the adaptability of the method to exploration of stress patterns carrying sharp gradients is well illustrated in Fig. 4, which shows the stress across and a little beyond a circular area of high yield strength in a soft steel plate. As shown, a 1 in. hole was filled with high strength weld metal, ground smooth and stress relieved, so that a (vertical) tensile stress could then be developed by bending -.When bent only a little, the elastic

Fig. 2

Geiger counter diffractometer used for strain determinations Institute of Technology).

(J. T. Norton, Massachusetts

50

Hatfield Memorial Lectures VoL II

Fig .. 3 X-ray determination of longitudinal stress distributions (curves A and B) in a notched plate for the two average stress levels indicated by horizontal broken lines G. T. Norton, Massachusetts Institute of Technology). strains were uniform across the specimen as shown by the crosses. When the bending was increased to cause a little yielding, the stress in the hard metal was quite high, and with greater bending the strong disc, now very highly stressed, picked up much of the load with consequent low stress in the plate material. Electronic methods of determining and recording X-ray intensities, combined with an extremely narrow slit and a suitable monochromatised beam, opens the way for the use of absorption to determine steep composition gradients as in diffusion.l" The actual

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Trends in Metallurgical Research in the United States

51

specimen is a thin section, cut diagonally if desired, containing the full diffusion zone. The specimen is moved in front of the very narrow slit and the intensity of the unabsorbed beam is continuously recordedas a function of the movement of the specimen. A typical result is shown in Fig. 5~in which a composition/distance curve is plotted for the full diffusion zone and compared with chemical analyses. In the study of phase diagrams of multiphase binary alloys this method is particularly effective in that, when phase equilibrium is established, an abrupt discontinuity may be revealed in the absorption curve corresponding to the composition limits of the individual phases, which in a single intermediate composition would exist together.

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Comparison of X-ray absorption and chemical methods for determining the nickel concentration in a Ni-NiAu diffusion zone: 720 h at 875°C.

Another technique for rapid determination of self-diffusion is that employed at the Massachusetts Instituite of Technology, 12 involving the use of a radioactive isotope and utilising the local radiation to darken a film which is appropriately brought in contact with a surface representative of the diffusion gradient. Figure 6 shows the result of micro densitometer measurement on a film applied to an oblique section across the interface zone.l ' Neutron diffraction is analogous in a general way to X-ray diffraction, but the mechanism is sufficiently different for the method to provide important advantages in certain types of crystal structure studies.!" The diffraction of neutrons by nuclei rather than orbital electrons is of value in studies of alloys of elements adjacent in the periodic table, as for example Ni3Mn and FeCo.15 The latter exhibits a greatly intensified diffraction line

52

Hatfield Memorial Lectures VoL II

0, tan a)t. sq.cm. x lOb

Fig. 6 Intensity / distance curve obtained by autoradiographic technique in study of self diffusion of gold in Au-Ni alloys (80 at.-%Ni) (M. Cohen, Massachusetts Institute of Technology). with neutrons for the superlattice line 100 as compared with the X-ray line, as shown by the following table: Relative Diffraction Line Intensities Characterising Order and Disorder in FeCo X-rays

1100 superlative line 1110 normal lattice line

1 1390

Neutrons

1 6

Because of the interaction of the spins of the neutrons and those of the 3D electrons of the atoms of the structures to cause a particular scattering (known as 'magnetic'), the magnetic structure of metals and alloys can now be penetratingly investigated by neutron diffraction. 16

ANALYSIS OF STEELS BY X-RAY FLUORESCENCE Another field in which the Geiger counter has brought about significant advances is X-ray spectroscopy or fluorescence analysis,"? the principles of which have been known for many years. Figure 7 shows plots of the intensity of the Ka line v. concentration for nickel, chromium and molybdenum in stainless steels as determined in our laboratories using the equipment shown in Fig. 8. These curves demonstrate that accurate analysis of stainless steels for these elements can be made by the X-ray spectrographic method. Recent developments in this field indicate that curved crystal analysers, which have definite advantages in resolution and intensity, will be widely used in the future, as will vacuum or helium filled optical paths to permit the determination of low atomic number elements. Birks and Brooks!" have recently described the use of a curved crystal spectrometer for the analysis of small concentrations of niobium, hafnium, tantalum, thorium and uranium in very small samples; also they report the use of a three channel, curved

Trends in Metallurgical Research in the United States 2000

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in

MARTENSITE IN SODIUM AT -400°F (-240°C) It will be recalled that Dr C. S. Barrett!" accomplished some very difficult experimentation in his discovery of a martensite-like transformation in lithium upon progressive cooling below about 70 K. Cold working, he found, even at somewhat higher temperatures, induces transformation to a face centred cubic structure, while the athermally transformed product was close packed hexagonal. Also in 1948, he reported a possible similar transformation in sodium at a somewhat lower temperature.s" The technique-" now employed in re-investigating the alkali metals uses the interesting apparatus of Fig. 9; in it accurate diffraction data (Cu radiation) are obtained with a

54

Hatfield Memorial Lectures Vol. II

Geiger or proportional counter, the specimen temperatures being carried down to about 1.2 K in the cryostat. No transformations were found in potassium, rubidium or cesium, but in sodium the transformation structure is clearly revealed as close packed hexagonal with imperfect stacking as in lithium.F-' Cold work merely produces more of the hexagonal modification, the axial ratio of which is 1.634 as compared with 1.637 for hexagonal lithium. The disparity is of interest in consideration of the theoretical effects of electron distribution.

Fig.9

X-ray diffraction apparatus used in studies of low temperature transformation in sodium (C. S. Barrett, Institute for the Study of Metals).

To obtain the X-ray diffraction data, the needed clean sodium surface is prepared in situ, after the pure helium atmsophere has been established, with an ingenious chisel manipulated by a 'sylphon' attachment, a tool adapted to scrape away oxide as well as to accomplish cold work when differently used. For crystallographic spacings and ratios, reflections from the atomic planes free of faults are used. The transformation first occurs upon progressively cooling below 36 K and may reach to 10% at 1.2 K. The metallographic evidence (Fig. 10) of the first small transforrnation-" is the geometrical roughening of a highly specular surface; the markings are preserved at room temperature. This 'polished' specimen (Fig. 11) is obtained by melting the sodium in a short glass tube into which it was vacuum distilled and sealed. The very smooth surface, resulting from fusion clear of the restraint of the glass, roughens with the first transformation which is clearly of the martensite type.

Trends in Metallurgical Research in the United States

Fig. 10 Transformation markings in sodium after cooling to 20.4 K (x2S) (C. S. Barrett, Institute for the Study of Metals).

55

Fig. 11 Sample of sodium in evacuated tube illustrating type of specimen used in metallographic studies of low temperature transformation (C. S. Barrett, Institute for the Study of Metals).

THE NATURE OF MARTENSITE Over a period of some five years a kind of controversy has been going on between two schools of thought about the nature of martensite transformation. Hollomon Fisher and Tumbu1124 have visualised a simple nucleation and growth mechanism forthe transformation. In their view, the effective nuclei, which at low temperature become supercritical in size, are located in solute poor regions. Growth is then by a diffusionless, atom by atom movement in the receding austenite as .the martensite advances through a coherent interface. The Cohen and Averbach-> group prefer to regard martensite nuclei as strain centres (strain embryos) wherein resides sufficient strain energy to initiate a co-operative displacement among the austenite atoms. If it occurs to any that this difference is something less than antipodal, be assured that the vigour of defence of the opinions gives no support to such a view. Some prefer their thermodynamics to be more mechanical, others more chemical. Supporting the 'strain nucleus' concept, and weakening the 'atom by atom shift' contention is the circumstance that no slowing of the transformation is observed at very low temperatures.ss while the existence and behaviours of isothermal martensite are construed to support the nucleation and growth view of Hollomon et al.27 With respect to isothermally transformed martensite, Shih, Averbach and Cohen-f find that its formation is greatly accelerated by the presence of some athermal or ordinary martensite. It forms by nucleation of new plates rather than by growth from the first plates within a grain. A strong argument for strain nuclei (embryos) is the observation that low temperature stress relief annealing seems, according to Machlin and Cohen.s? to destroy such nuclei while cold work produces them. As may be inferred from Fig. 12, Patel and Cohen-'? verified experimentally that the Ms temperature is raised by uniaxial compression and evenmore by tension, and lowered

56

Hatfield Memorial Lectures VoL II

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by hydrostatic pressure, just as predicted by the sums of the chemical free energy changes and the work done on or by the external forces. Another aspect of the martensite transformation that has been the subject of intense study is that of the formulation of specific atomic displacements, consistent with the observed crystallographic relationships between the parent austenite and the transformation product. An important contribution was made by Geisler.P! who pointed out that the irrational features of martensitic structures might be regarded as the result of plastic flow accompanying the transformation, and that the dominant role of the shear is to minimise the strain energy and not necessarily to generate the new crystal structure. Perhaps the most successful detailed analysis is that developed by Wechsler, Lieberman and Read.V in which the transformation is visualised as involving a pure distortion somewhat of the type in Fig. 13, suggested by the present speaker.P combined with a rigid body rotation. On this basis a theoretical analysis is developed which predicts the habit plane, orientation relationships and macroscopic distortions from a knowledge only of the crystal structures of the initial and final phases.

TEMPERING The current views from all quarters with respect to the mechanisms of tempering are becoming more and more reconcilable and, indeed, there is something approaching unanimity on broad lines.P? Quite generally, for convenience, tempering of quenched steel is regarded as occurring in four distinct, though not necessarily independent, stages.t> (I) The first stage of tempering high carbon martensite is regarded as embracing the rej ection of the first supercritical particles of carbide, and the changing in degree of

Trends in Metallurgical Research in the United States 57

(0)

(e)

Fig. 13

The distortion which any homogeneously transforming volume of austenite undergoes in order to become martensite.

martensitic tetragonality corresponding to a much lower carboncontent, as proposed by Jack.36 Roberts, Averbach and Cohen-'? confirm this view; in some very precise work, they estimate the carbon content of the hexagonal epsilon carbide as corresponding to a hypothetical stoichiometric Fe2.4C and the new impoverished martensite as of 0.25% carbon content. On the basis of observations with the electron microscope, Cohen suggests that the martensite in hardened steel inherits a network of sub-boundaries from the parent austenite, and that these locations provide the sites for the first carbide precipitation upon tempering, the carbide being visible as a network around the sub-grains as shown in Fig. 14a.

Fig. 14 Electron micrographs (x1S,OOO) of 1.4% carbon steel quenched from 2200°F and tempered 1 h in the range 4S0-600°F: (left) 450°F; (middle) SOQoF; (right) 600°F (B. S. Lement, B. L. Averbach and M. Cohen, Massachusetts Institute of Technology).

58

Hatfield Memorial Lectures VoL II

(II) The second stage, in carbon steels at least, is that of an isothermal transformation of any residual austenite, the rate of which may be governed by the diffusivity of carbon in austenite. (III) The third stage is characterised by the formation of cementite of regular Fe3 C crystal form, through re-solution of the hexagonal carbide and re-precipitation along the martensite plate boundaries and within the martensite itself Cohen et al.38 show an early stage of the formation of cementite in Fig. 14a, while in Fig. 14b the cementite has grown into fairly massive networks and the fine subboundary films of epsilon carbide are being redissolved. In Fig. 14c, the epsilon carbide has completely disappeared in favour of the more stable cementite. As the tempering temperature is raised further, the cementite networks along the boundaries of the martensitic plates thicken and become more continuous. This is accompanied by a resolution of the cementite particles that previously precepitated within the martensite. By this time, the matrix has lost its tetragonality and is indistinguishable from body centred cubic ferrite. The ferrite is outlined by the cementite networks, and inherits an acicular configuration from the prior martensite. Eventually, the cementite undergoes coalescence and the well known spheroidised structure is attained. There is a concomitant migration of the ferrite grain boundaries, resulting in grain growth and equiaxed shapes. According to Cohen.P? the so called 50QoF (260°C) embrittlement sets in when the cementite begins to form more or less continuous networks around the boundaries of the martensitic plates. Retained austenite, if present, also transforms in this temperature range into a bainitic product, but this is manifestly not the cause of the 500°F (260°C) embrittlement because the latter is not reduced in magnitude when the tempering is carried out on refrigerated samples, containing less retained austenite than as hardened samples. The 500°F (260°C) embrittlement becomes an important consideration in the very high strength steels that have to be tempered in the low range of 400-500°F (204260°C), for maximum useful strength. Kinetic studies.f" however, have shown that silicon is very effective in retarding the nucleation and growth of cementite. Possibly the reason for this striking phenomenon is that the cementite does not accept silicon in its lattice to any appreciable degree, and therefore this stage of tempering is controlled by the rate of silicon diffusion away from the cementite rather than by the rate of carbon diffusion towards it. Be that as it may, the presence of about 1.5% silicon raises the embrittling range by some 200°F (110°C), and hence such steels can be tempered successfully at 400-S00°F (204-260°C). At the same time, there is a noticeable strengthening effect from the silicon in solid solution. (IV) This interesting stage is prominently manifest in steel containing alloying elements which fonn special carbides. Otherwise only coarse spheroidisation from fine platelets and particles takes place. In the presence of chromium, molybdenum, tungsten or vanadium, a new double carbide, characteristic of the composition, fonns from new nuclei, while the cementite discs or platelets disappear through solution.

Trends in Metallurgical Research in the United States

59

This phenomenon has never been so well shown, so far as I am aware, as in the electron micrographs obtained by the associates of Professor Quarrell+! at the University of Sheffield. It is a matter of great interest to the present speaker to see today the actual photographic evidence supporting a theory42 offered thirty years ago by him with scarcely enough evidence, but with a conviction that no other mechanism could then so well explain the phenomenon of secondary hardness.

THE BORON EFFECT The discovery of the boron effect on the hardenability of steel was made at a time when the ordinary alloying elements added to enhance hardenability were in short supply.F' and it has served very well. It had been amply shown and reverified that the only known way to realise toughness and high strength simultaneously was to replace the fine pearlite microstructure of unhardened steel with the tempered martensite (or lower bainite) structure. One can, for example, realise 130 tons/In.s at ordinary temperature, coupled with an elongation in 2 in. of better than 8%, but only in hardened and tempered steel. Still greater plasticity can, of course, readily be coupled with somewhat lower strength. The practical problem then was simply to obtain and add alloy enough to ensure the needed depth of hardening in the needed tons of large sections! Boron frequently almost doubled the diameter of rounds which could just be hardened through, and as little as about 0.0015% accomplished this (Fig. 15). How it accomplished the amazing retardation of transformation to fine pearlite or upper bainite was, and is, obscure. It has attracted considerable research in our own laboratories, and as a result, we are somewhat acquainted+' with its effects. We know, for example, that it serves best with lower carbon steels, and that it can be made ineffective by heat treatment. But the mechanism of its influence is not yet known. In my opinion the most important observation has been made by Nicholson+> at the University of Chicago, and shown in Fig. 16. He has found that in essentially pure, dilute, iron-boron alloys the depression of the temperature of transformation with increasing cooling rate is steeper than for the boron free iron. Considered along with the greater effect in lower carbon steels it would appear that boron may somehow 'denucleate' ferrite. Some ten years ago the present speaker noted that the boron effect was not far from that exerted by an austenite coarsening of two or three ASTM grain size numbers, and indeed that boron steels were usually a grain size number or two coarser at quenching temperature, and somewhat facetiously suggested that perhaps boron merely ensured inactive, non-nucleative, glassy inclusions instead of the active crystalline, aluminous inclusions, ordinarily found in our alloy steels. It may probably be much more complex than this. Nicholson+" has found also that 80% of the carbon in Fe3C can be replaced by boron with a change in the b lattice parameter from 5.07 to 5.40 A and a rise in the Curie temperature from 374 to 1040°F (190 to 560°C) and with an increase in saturation magnetisation. These data, however, do not seem at present to suggest a definite boron mechanism.

60

Hatfield Memorial Lectures VoL II 4·0

3~

k--

3·5

<[ _..•••

/

>~

~

--

~

- IL

~ ~ 3·0 % u 0« -,

2·5 c

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V

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~

J'\.

1-0

0

0-0004

0-0008

BORON ADDITION,

Fig. 15

0·0012 0/0

Effect of boron additions on hardenability.

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-," •. ···~·

laco

/NRciron

~: ..::.::::::'.. - : : . ~~?5%NH.

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olio)

. _.I_j;:.:.ta..... • I alloy

~

••.

~ I

I

bOO o'-------Io...l.o~O~2~OO'=-O~3~O~O~O--:-40~O~O::-----:5~O:""='O-=-O---:--::bO~OO COOLING VELOCITY, °C./s

Fig.16 Effect of cooling rate on 'Y--t a transformation temperature in pure iron, O.5%Ni-Fe alloy and O.007%B-Fe alloy: 950°C austenitising temperature (M. E. Nicholson, Institute for the Study of Metals).

ELECTRON METALLOGRAPHY In the field of metallography, electron microscopy and its companion tool electron diffraction have been combined in a powerful technique by R. M. Pisher,"? of our laboratories. He isolates, in a thin plastic film or replica, minute particles of a segregate phase, so that they maintain in the replica the same relative position as they occupied in the metal matrix. This method is known as the extraction replica technique. The principle of the method is shown schematically in Fig. 17. A conventional plastic replica is applied to the metallographically polished and etched surface of the sample, and then the surface is re-etched through the replica, which is quite permeable to etching

Trends in Metallurgical Research in the United States

61

FIRST ETCH

SECOND eTCH

~ Fig. 17

REPLICA

Schematic diagram of extraction replica technique.

solutions. The effect is twofold: first, to provide a conventional replica of the topography of the etched surface, and second, to undermine and extract particles of the segregate phase so that they will adhere to the replica upon stripping. By this means, the actual particles may be viewed directly at very high magnification in the electron microscope to determine their size and shape and their orientation and distribution in the matrix. Furthermore, the method provides an ideal arrangement for mounting the particles for determination of their crystallographic identity by electron diffraction. The use of this new technique, in combination with selected area electron diffraction, also makes it possible for the first time to isolate individual minute particles of a solid phase precipitate, and to obtain crystallographic information from a single particle that has been previously observed and selected for study under high magnification in the electron microscope. One illustration of the use of this technique is the observation and identification of precipitate particles in the grain boundaries of a low carbon (0.026%) 18/8 Cr-Ni steel (type 304L) , which had, notwithstanding, developed grain boundary susceptibility to corrosive attack. The precipitate particles are shown in the electron micrograph of an extraction replica (Fig. 18) at a magnification of xts, 000. The particles are roughly of triangular shape and measure about 7000 A on a side. Electron diffraction patterns of the particles in this replica demonstrate conclusively that they are the complex ironchromium carbide (Fe,Cr)23C6' For the most part they are situated within a rapid etching zone at the grain boundary. This rapid etching zone undoubtedly corresponds to the grain boundary region which is depleted of chromium, and X-ray diffraction evidence indicates the presence of ferrite in this thin zone.

62

Hatfield Memorial Lectures VoL II

Fig. 18

Electron micrograph (x15,OOO) of grain boundary in 18/8 low carbon steel, showing carbide precipitate and zone of chromium depletion.

Another interesting illustration of the use of this technique is the observation and identification of the constituent that precipitates during ageing of a copper bearing steeL The rod-like particles and the electron diffraction patterns identifying them as copper are shown in Figs. 19 and 20 for ageing at 1200°F (649°C) and 1300°F (704°C). It is also pertinent to mention here the work of Subcommittee XI of ASTM Committee E-4 on Metallography'P" on the morphology of the bainitic and tempered martensitic structures in eutectoid carbon steel. One interesting observation is that upper bainite, in an early stage of austenite decomposition, consists of ferrite needles with stringers or chunky particles of cementite at the needle boundaries. These needles, which appear to be lath-like, nucleate at austenite grain boundaries and grow into a grain of austenite in a

Fig. 19 Electron micrograph of copper particles precipitated during ageing of a 1.1 %Cu steel aged for 72 h at 1200°F.

Fig. 20 Electron micrograph of copper particles precipitated during ageing of a 1.1%Cu steel aged for 72 h at 1300°F.

Trends in Metallurgical Research in the United States

63

parallel array. This situation is shown in Fig. 21. This sample of eutectoid carbon steel was 20% transformed at 8S0oP (454°C). In the latter stage of transformation, the remaining austenite decomposes to fine pearlite. The mixture of upper bainite and fine pearlite in a eutectoid carbon steel fully transformed at 9S0 P (510°C) is shown in Fig. 22. These observations on upper bainite tend to support the mechanism of formation suggested by Hultgrenr'" it may be considered to be very similar to that of the very earliest stage of formation of a pearlite colony. Another interesting finding is that lower bainite consists of ferrite needles containing within their boundaries very fine platelets of carbide that are generally oriented at an angle of about 55° to the needle axes, as shown in Fig. 23. This sample of eutectoid carbon steel was completely transformed to bainite at 500°F (260°C). This observation, of course, suggeststhat the carbide platelets are precicipated from ferrite supersaturated with carbon almost concurrently with the formation of the ferrite from austenite. Thus it appears that the mechanism of formation of lower bainite is similar in some respects to the formation of martensite. Looked at in this way, the mechanisms of formation of upper bainite and of lower bainite seem to form a logical transition between the mechanism of formation of pearlite and that of martensite. This brief discussion of high temperature solid phase transformations brings to mind the perennial hope that in the not-too-distant future it may be possible to observe directly the progress of transformations, such as the pearlite and bainite reactions, rather than to be forced to rely on 'post-mortem' examination of quenched structures. Some years ago, Rathenau and Baas"? of The Netherlands, and others, indicated the potentialities of thermionic emission microscopy for such studies. More recently, Heidenreich=? of the 0

Fig. 21 Electron micrograph (X15,OOO) of upper bainite in a eutectoid carbon steel (par-

tial transformation at 850°F).

Fig. 22 Electron micrograph (x15,OOO) of upper bainite and fine pearlite in a eutectoid

carbon steel (complete transformation at 950°F).

64

Fig.

Hatfield Memorial Lectures V01. II

23

Electron micrograph (x15,OOO) of lower bainite formed at 500°F in eutectoid carbon steel.

Bell Telephone Laboratories has constructed an electrostatic thermionic emission microscope and has used it to study the transformation of austenite to ferrite and pearlite in iron and iron-carbon alloys. Figure 24 shows a thermionic emission image of pearlite taken at 1292°F (700°C) of a O.77%C iron-carbon alloy. It illustrates the resolution and contrast that presently can be obtained by this technique. The series of emission micrographs of a

10 J.L Fig. 24

Thermionic emission micrograph (X3500) taken at 700°C of pearlite in a 0.77% carbon Fe-C alloy (R. D. Heidenreich, Bell Telephone Laboratories).

Trends in Metallurgical Research in the United States

65

0.22% carbon steel at 1490°F (8100e), shown in Fig. 25 demonstrates the use of this technique to follow the growth of a ferrite grain from austenite. It is reasonable to expect that in the coming years the technique will be developed to the point where image interpretation is easier and more certain, and that improvements will broaden the range of temperature over which it can be applied.

Fig. 25 Thermionic emission micrographs (x850) showing the growth of a ferrite grain from austenite in a 0.22% carbon steel at 8l0oe (R. D. Heidenreich, Bell Telephone Laboratories).

A MECHANISM OF DIFFUSION Along with accelerated studies of diffusion, particularly self-diffusion, brought about by the availability of numerous radioactive isotopes, renewed interest is taken also in more accurate formulations and mechanisms of diffusion. Particularly engaging is Nachtrieb's so called 'relaxed vacancy' model 51 for diffusion. He notes that hitherto the value of the vacancy mechanism has been more conceptual than quantitative. Nachtrieb prefers to consider the creation and movement of a 12-14 atom region of disorder relaxed around a vacancy rather than the single vacancy alone; such a region would, he reasons, have an energy content about equal to that of a similar number of atoms in the liquid state. It is consistent with such a theory that the activation ene:rgy for self-diffusion in cubic metals is

16.5 times the latent heat of fusion, as may be seen in the following table:

66

Hatfield Memorial Lectures Vol. II Metal

Latent Heat of Fusion, cal./mole

Activation Energy for Self-Diffusion, cal./mole Observed Calculated

Au Ag

3060 2730

Cu

3110

{

Pb

1190

{

Na Co Fe

636.2 3700 3630

51,000 45,950 47,140 54,000 27,900 24,120 10,450 61,900 59,700

50,280 44,850 50,920 19,550 10,450 60,790 59,640

Further support is given by observations of the dependence upon pressure of the rate of self-diffusion in sodium (Fig. 26). By taking into account the higher melting point associated with higher pressure, the rates of diffusion can be brought into agreement. Diffusion behaviour of silver in silver base alloys wherein the solute lowers the melting point are in agreement with Nachtrieb's ideas. A critical experiment would be the selfdiffusion in silver-palladium alloys (rising liquidus), since here a lowering rather than an increase in the self-diffusion rate would be predicted for increased solute concentration.

-,

t\.

"-CJ\.

~

\

'"

-,

'"'""- '\.

~_"\..

-,

\

H= 12,ObO col./mole p = 8000 kq./sq.cm.

H = 10,450 col./mole_ P I kq./sq.cm.

=

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Fig. 26

Effect of externally applied pressure on the rate of self-diffusion in sodium (N. H. Nachtrieb, Institute for the Study of Metals).

DISLOCATIONS Rapid progress has recently been made in the refinement and usefulness of basic dislocation theory through the imaginative and creative concepts of Cottrell.V Frank53 and

Trends in Metallurgical Research in the United States 67 others in the United Kingdom, and supported by Read.>' Shockley-> and others in the USA. Striking confirmation of the nature of dislocations has been provided by the agreement between a number of experimental observations and predictions based on dislocation theory. If the dislocation model for a low angle crystal boundary is correct, it should be possible to cause such a boundary to move under the action of an applied stress. Parker and Washburn 56 have shown that such movements do indeed occur. A specimen containing a low angle boundary was loaded as a cantilever beam, as shown in Fig. 27. The resulting boundary movement is illustrated in Fig. 28, in which reversal in motion produced by a reversal in stress is also shown. Systematic studies of the effect of low angle boundaries on the behaviour of crystals subjected to stress are providing further insight into the mechanisms of strain hardening and creep. 57

Fig. 27 Cantilever beam specimen used in studies of the movement of low angle crystal boundaries under applied stress (E. R. Parker and J. Washburn, University of California).

Fig. 28 Photomicrographs (x50) illustrating movement of a low angle crystal boundary under applied stress; the boundary first moves to the right and then to the left upon reversal of stress (E. R. Parker arid J. Washburn, University of California).

68

Hatfield Memorial Lectures VoL II

Another confirmation of the dislocation model has been provided by observations of the spacing of conical etch pits along low angle boundaries. This approach is perhaps best illustrated by the work at Bell Telephone Laboratories on crystals of germanium.f" According to the dislocation model, the spacing of dislocations will depend on the mismatch in orientation at the boundary and can be calculated. Striking confirmation of this concept is shown in Fig. 29, wherein the spacings of etch pits along four intersecting boundaries are clearly different. The agreement between predicted dislocation spacing and observed etch pit spacing for various degrees of orientation difference is shown in Fig. 30. From such observations, it can be reasonably concluded that the centres of the etch pits mark the points of emergence of edge dislocations.

Fig. 29 Photomicrographs (x4S0) showing intersecting low angle lineage boundaries in germanium single crystal CW. G. Pfann and L. C. Lovell, Bell Telephone Laboratories). Observation of the disposition of dislocation etch pits after various straining and annealing treatments is proving a fruitful approach to a better understanding of the role of dislocations in slip, recovery and recrystallisation processes. Gilman''? recently observed an array of etch pits along a basal slip plane in a zinc crystal (Fig. 31) that is believed to have resulted from the 'piling-up' of dislocations at an obstacle during slip. Dunn and Hibbard''? are investigating the nature of polygonisation in 3Y4% silicon ferrite by observations of dislocation etch pits. The series of photomicrographs shown in Fig. 32 illustrates the changes occurring during annealing of a bent single crystal. The early changes shown are interpreted as reflecting glide and climb of edge dislocations, the changes presumably reducing strain energy.

SPIRAL GROWTH OF CRYSTALS Since the role of screw dislocations in the mechanism of crystal growth was first proposed by Frank, observations of growth spirals as predicted by the theory have been made on a

Trends in Metallurgical Research in the United States

69

b

\0 '¢ I

Q x

5 o

4

Q Z

\ ~

U ~

~bSeryed

0.. V)

Z

o

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u

9

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=

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e

r-,

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4 ANGLE OF TILT 6, rediens x 10-4

Fig. 30 Comparison of computed dislocation spacings and observed etch pit spacings at low angle crystal boundaries as a function of differences in orientation (F. L. Vogel, Bell Telephone Laboratories).

Fig. 31

Etch pits lying along a basal slip plane in a zinc crystal (x600) Electric Company).

a.

Gilman, General

70

Hatfield Memorial Lectures VoL II

(a)

(b)

(c)

Fig. 32 Etch pits in a single crystal of a 3.25% silicon ferrite bent round [211] axis and observed on (211) plane (X750): (a) as bent; (b) 1 h at 650°C; (c) 1 h at 700°C; (d) 1 h at 850°C; (e) 1 h at 950°C; (f) 1 h at 1050°C (C. G. Dunn and W. R. Hibbard, General Electric Company).

wide variety of crystal surfaces. Recently Pollock and Mehl=! have observed spiral formations during the growth of cadmium crystalsfrom vapour at controlled supersaturations in a special growth cell. The observed spirals, visible at x75, are grown on a cadmium substrate, formed by subliming cadmium on a cold surface; spiral growth occurs, apparently, only when the condensed substrate is melted and frozen. The spirals, which have step heights between 300 and 1500 A, as determined by interferometric measurements, form at ridge boundaries, i.e. boundaries separating two regions that differ slightly in orientation as evidenced by a small angle between their basal surfaces. Several stages in the formation of a spiral are shown in Fig. 33. An interesting pattern that may be in some way related to the crystalgrowth mechanism was recently observed in a copper crystalby Simnad=?of the Carnegie Institute of Techno logy (Fig. 34). The pattern was developed on the (111) face of a single crystal of copper annealed at 1832°F (1000°C) for 12 h in argon. No similar pattern was observed on the (100) or (110) faces. FILAMENTARY CRYSTALS Considerable interest has centred recently in. various forms of filamentary crystalline growth, frequently called 'whiskers,' which were first observed as a spontaneous growth

Trends in Metallurgical Research in the United States 71

Fig.33

Photomicrographs (X7S) illustrating spiral crystal growth in cadmium (W. 1. Pollock and R. F. MeW, Carnegie Institute of Technology).

Fig. 34

Photomicrograph

(xSOO) of thermally etched (111) surface of copper single crystal (M. Simnad, Carnegie Institute of Technology).

on tin, zinc and cadmium plated components in commercial telephone equipment. 63 These filaments were found to be single crystals of the metal, generally about 2 J..lmin diameter, although a considerable range of sizes has been observed.v? The spontaneous linear growth rate is of the order of 1 mm per year, but it has been found that application of pressure will cause a growth of 1 or 2 mm in length in a few hours.v" for crystals of this type grow from the base,66 as shown in Fig. 35. Whiskers of iron, silicon and copper formed by condensation from a vapour phase by depositon at the tip, rather than from the base."? Filamentary crystals of oxides and sulphides have also been grown under special conditions. 68 Regardless of how they are grown, all these filamentary crystals have unique mechanical properties.v? While under microscopic observation they have been strained elastically as much as 1.5% (Fig. 36), indicating an elastic limit 100 times as great as that of large

72

Hatfield Memorial Lectures VoL II

Fig. 35

Electron microscope silhouettes (x22S0) showing stages in the growth of a tin whisker (S. Eloise Koonce and S. M. Arnold, Bell Telephone Laboratories).

Fig. 36 Bending and recovery of tin whiskers: (top) radius of curvature 0.009 ern; (bottom) same whisker after removal of constraining probe (C. Herring and J. K. Galt, Bell Telephone Laboratories).

single crystals. Similarly, their creep rate at loads exceeding the yield strength of the bulk metal has been found to be many times smaller than the usual value. Bent beyond their elastic limit they develop a sharp bend, as shown in Fig. 37. Iron whiskers of approximately 0.001 in. diameter have exhibited a yield strength of 270,000 lb in.? for the outer surface (in bending) and a yield strength of85,000 lb in.2 in microscopic tensile tests.?" These values are to be compared with yield strength of 4000 lb in.2 for conventional single crystals of iron. The ex ~ 'Y transformation temperature is reported to be about 2012°F (1100°C), which is nearly 360°F (200°C) above that for the bulk metaL71

Trends in Metallurgical Research in the United States

Fig. 37

73

Typical permanent bend in tin whisker produced by stress in excess of elastic limit (C. Herring and]. K. Galt, Bell Telephone Laboratories).

The unusual strength of the filamentary crystals, approaching the theoretical value for a perfect crystal, is considered to indicate that they are nearly free of dislocations, and it has been suggested that the slender crystals might serve as seeds upon which to grow exceedingly strong macroscopic filaments.

METALLOGRAPHICASPECTS

OF FATIGUE BEHAVIOUR

A novel approach to the problem of the fundamental aspects of fatigue behaviour of the light metal alloys is being used by the Aluminium Research Laboratories of the Aluminium Company of America utilising metallographic techniques.V The testing of fatigue specimens which have been completely chemically polished is interrupted periodically in order to obtain plastic replicas of the surface. These replicas then provide a permanent record of the deformations occurring in particular areas throughout the test. They may be examined at all magnifications with the optical microscope and, using further techniques, with the electron microscope. All stages of fatigue deformation are observable with the replica technique from the innocuous stages involving slip to the more insidious stages involving cracking (Fig. 38). The formation and propagation of very small fatigue cracks has been observed and studied, especially as to how they relate to such microstructural features as constituents and grain boundaries. Success has been achieved in relating the number of cycles at which each new fatigue induced stage is entered to the number of cycles needed to cause failure. Points representing the first appearance of slip in the specimen, the first cracking, etc., fall on smooth

74

Hatfield Memorial Lectures VoL II

Fig.38 Photomicrograph (x250) of plastic replica taken from surface of an Al-G.97%Mg alloy during a fatigue test. Slip saturation and crackjoining after 148,700 cycles at a stress of6250 Ib/in.? are shown (M. S. Hunter and W. G. Fricke, Jnr, Aluminium Company of America). curves having the same general shape as the conventional SIN curve denoting failure, as shown in Fig. 39. Quantitative measurements of the amount of deformation present at any time in the tests are obtained from the replicas using a grid intercept method. The influence of the factors of stress, number of cycles, grain size and alloy content on the amount of deformation is being evaluated further by the metallographic method. The applicability of the technique to ferrous and other metals is being developed .

~

. cT ~

-~.~r--t-=9--~-t----+-----+----I

6500~--+---3or

5500~

~

__

~

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10

Fig. 39 Progression of fatigue damage in an Al-0.97%Mg alloy as revealed by microscopic examination of plastic replicas of specimen surface at various stages of cycling (M. S. Hunter and W. G. Fricke, Jnr, Aluminium Company of America).

MACHINABILITY A great deal of attention is being focused on the fundamental factors influencing the ease with which a metal can be cut, i.e. machinability, especially from the point of view of

Trends in Metallurgical Research in the United States 75 isolating the basic material characteristics responsible for good machinability. Studies of force relationships and chip formation during machining, together with theoretical analysis developed by Piispanen.Z'' Merchant.?? Shaw 75 and others constitute one important line of attack. As an example, these factors were studied in a series of three low carbon steels containing sulphur in amounts varying from 0.025 to 0.250%, the sulphur additions being made to successive ingots of an individual heat.?" It was found that the superior machinability of the higher sulphur steels accompanies decreased friction between chip and tool and reduction in shear strain during chip formation. The decrease in friction and shear strain reduce the work required for removal of a unit volume of metal during cutting. The effect of increased sulphur content on friction force for various feeds is shown in Fig. 40.

/-

180

/

n

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140 ~I

~ 120

2 z

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100

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80 I

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40

Fig. 40

0-001 0-002 0-003 0-004 0-005 0-006 FEED. in.Zr ev,

Effect of sulphur content on the friction force developed during machining at 202 surface ft/min in a constant pressure lathe; 0.14%C, 0.79%Mn alloy.

Chao and Trigger"? have contributed greatly in the direction of the relation between chip formation and tool life through their theoretical considerations of temperature distribution in the cutting tool. They have pointed out that the tendency for 'welding,' and even diffusion, to occur at the tool/chip interface may play an important role in tool life.

DEFORMATION AND BRITTLE FRACTURE Few, if any, subjects in the physical rnetallurgy of iron and steel have been accorded in America the overall research man hours and effort that have been given to the problem of brittle fracture. The modern viewpoint emphasising the actual local state of stress in considering flow in metals had been already fruitful; nevertheless, the investigations were

76

Hatfield Memorial

Lectures VoL II

accelerated and increased greatly at the time of the numerous fractures in ship hulls, at that time inexplicable. These investigations have quite naturally fallen into two categories: (i)

The development of tests which should evaluate the suitability of a steel, usually as welded, for designated structural purposes (ii) The theoretical and experimental inquiries having as objectives greater understanding of the mechanism of brittle fracture and the factors and influences which bring it about.

These two major types of approach were, quite clearly, mutually helpful, and have adduced tremendous advances in the knowledge of deformation and fracture in metals. Generally speaking, the search for mechanisms which deal with actual individual grain behaviours have been more rewarding in America than the statistical or continuum theories. The thinking has been greatly influenced by the gamut of splendid observations by Robertson.?" Williams,"? and Pellini, Puzak, et aI.80 It has seemed obligatory to weigh heavily the good correlation of transition temperatures of the ship plates of known service performance as determined for an energy level of 15 ft lb, Charpy V notch, and as determined by the Pellini crack starter, dropweight test, and the crack-starter, explosion test. The latter confirm the findings illustrated in the following table Service Temperature Charpy V notch Energy, ft lb 3.2-11.7 3.5-19.0 9.2-51.4

(7.6 av.) (9.3 av.) (19.2 av.)

Average 15 ft lb Transition Temperature

Plate Service Performance

900P (32°C) 68°F (20°C) 39°F (4°C)

Fracture started Fracture through Fracture stopped

Some of the generally acceptable conclusions are as follows: (1) Sharp notches, such as cracks, freshly started in some cases, are an appropriate fundamental condition for experimentation. Misadventure in building structures, particularly in welding, may produce the equivalent of sharp notches or cracks; hence special interest centres in the ability of the metal in an element of structure to resist the advancement of a crack, under certain definable conditions of temperature and stress. (2) Transition temperature, with a precisely designated criterion, is a useful item of information. In practical testing, the degree of deformation prior to rupture, inherent in the criterion of transition temperature, should reflect the conditions of intended service. The amount of deformation selected as the criterion in the variety of tests has covered the range fromalmost none to a very great deal, and has been observed through reduction in area, fracture texture and energy absorption. This determining criterion of transition temperature (be it a limiting deformation, an energy value or a fracture texture) should be

Trends in Metallurgical Research in the United States

77

in keeping with the condition of loading which the structure is intended to meet. If, through a very little flow, stress ahead of a notch is substantially reduced, as in the case of high residual stress, then only a little plasticity under notch conditions is needed to prevent crack propagation. If this stress persists even though considerable flow occurs, then a more plastic material (lower transition temperature, generally) will be required; i.e. one which, at service temperature, cracks only with the accompaniment of considerable deformation. (3) The rate of external loading is not of as great influence, especially when the notch is sharp, as was once supposed. It is reasoned that the actual rate of increase in stress and hence in any resulting strain is, in any event, very high just ahead of the crack. (4) The test employed should as nearly as possible approach in temperature and notch condition those contemplated in the proposed service which the material is designed to meet. To some extent, a temperature difference may compensate for a change in notch (triaxiality of tension) or rate of loading, but such extrapolations should be used with great caution. (5) By way of a mechanism of brittle fracture, the following is envisioned. Few indeed are the metals, even in the state of stress just ahead of a crack, which rupture with no plastic deformation whatsoever, prior to the break. In steel, even in the so called 'nil ductility' temperature range, some few grains at least are so oriented as to incur slip (deformation), and in so doing, pass on the load to other grains unable to flow. Some of these grains, under high disruptive stress, crack well ahead of the principal or gross crack (Fig. 41). Hereupon, some grains again flow, but some rupture and in so doing they connect up smaller cracks and ultimately join with the principal crack. A little

Fig. 41

Sketch illustrating advancing cleavage crack, showing cracking ahead of main crack and various ways in which plastic flow may accompany an essentially brittle fracture.

78

Hatfield Memorial Lectures VoL II

lower temperature would (1) decrease the total of plastic deformation and hence (2) decrease the energy involved in rupture, and (3) increase the total of purely cleavage grain ruptures. A slightly higher temperature would (1) increase the proportion of grains undergoing slip, and also the total deformation, and hence (2) the energy involved, and (3) decrease the proportion of' cleavage' rupture in grains. At some point of decreasing ratio as between, broadly, 'brittle' and 'ductile,' the metal would just be able to halt a crack, or fail to propagate one. The degree of plasticity required for this manifestation depends upon the nature of the loading in respect to its maintenance of stress at the crack. Or, as stated by Irwin.f" spontaneous crack propagation will not occur until the strain energy released by an incremental extension of the crack exceeds the work absorbed by the material as a result of this extension. Supporting the endeavours to find criteria of the notch resistance of steels are a number of investigations dealing with microscopic scalephenomena, such as the work of Dr J. R. Low. That even the most brittle fractures in metals are initiated by plastic flow now appears to be well established by these approaches. Since plastic flow by slip involves the generation and motion of dislocations in crystalline materials, a number of models of processes by which dislocation motion might lead to microscopic and sub-microscopic cracks in metals have been investigated, mainly theoretically. It now appears probable that once any slip has occurred, a number of small cracks might be generated. Whether or not the crystal or the aggregate of crystals fails brittley before any further flow can take place then depends on the factors governing the growth of these submicroscopic cracks. Low, who so effectively correlated grain size with fracture stress.f? has been studying the surfaces of cleavage cracks in single crystals and polycrystalline aggregates to learn more about this crack propagation process.s> He finds that these cleavages are never perfect and that the cleavage surface always contains steps. These steps represent lines of overlap of the crack travelling on two slightly different levels and further tend to slow the crack propagation because of the energy absorbed in plastic deformation of the material in the region of overlap, which must be tom apart. Cleavage steps, according to Low, may arise in a number of ways: (i) They may originate at dislocations already present in the crystal (ii) They may originate at low-angle boundaries, as shown in Fig. 42, which is the cleavage surface of a single crystal of 3.25% silicon ferrite that contains a low-angle twist boundary (angle of twist: 1 or 2°) (iii) If plastic flow can occur at the tip of the crack, a high density of such steps is produced, as shown in Fig. 43. Here, a high velocity crack was stopped and then immediately started again at liquid nitrogen temperature; all along the crack front when it started again there is a high density of cleavage steps. This is interpreted as follows: so long as the crack was moving rapidly, there was not time for flow to occur ahead of the crack. Once stopped, however, flow could occur in starting the crack again, and this plastic flow, just ahead of the crack, caused imperfections which lead to the high density of cleavage steps shown.

Trends in Metallurgical Research in the United States 79

Fig. 42

Cleavage surface of a single crystal of 3.25% silicon ferrite which contains a low angle twist boundary (xl00) G. R. Low, General Electric Company). Crack stopped and started again

!

Direction of cracking -

Fig.43 Cleavage surface of a single crystal of3.25% silicon ferrite in which a high velocity crack was stopped and then started again, thus producing a high density of cleavage steps G. R. Low, General Electric Company).

80

Hatfield Memorial Lectures VoL II

At Columbia University, Dr Max Gensamer has been studying the mechanical behaviour of iron and steel at low temperatures, directed toward understanding low temperature brittleness. At temperatures in the neighbourhood of the boiling point of nitrogen at atmospheric pressure, the stress/strain curve of low carbon steel in simple tension is quite temperature dependent, the stress required at a given strain rising as the temperature is lowered, with a change in shape of the stress/strain curve. This is a temperature range in which carbon and nitrogen diffuse very slowly, so that if the diffusion of an interstitial impurity is involved in strain hardening, it would seem that the interstitial impurity involved must be hydrogen. To demonstrate that hydrogen does diffuse rapidly at such low temperatures he has studied the variation with temperature of internal friction and has reporteds+ an internal friction peak at about 120 K in steel charged with hydrogen (Fig. 44). He has since modified his apparatus to work at very low temperatures, using liquid helium, and has found85 another and more pronounced peak at about 50 K. This lower temperature peak may be associated with the strain induced diffusion of hydrogen in the iron lattice, while the higher temperature peak may possibly be associated with the movement of dislocations. The possibility of interstitial diffusion at these temperatures suggests observable ageing phenomena analogous to those attributable to carbon and nitrogen in steel at and above room temperature.

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Fig.44 Internal friction v. temperature for a 0.20%C steel: (a) annealed at 975 K for 40 h, 10-5 mm Hg; (b) pickled in HCl; (c) re-annealed at 775 K for 4 h, 5 X 10-8 mm Hg (L. C. Chang and M. Gensamer, Columbia University).

Trends in Metallurgical Research in the United States 81 ZONE MELTING We should be very remiss were we to omit from our discussion of trends the much publicised subject of zone melting for purification. It is applicable to any solvent-solute system in which an appreciable difference in solubility exists as between the molten and solid states. The principle is not wholly new. In making a single crystal by slowly withdrawing a solid bar from a melt a degree of the effect has resulted; even in ingot freezing, wherein the centre segregation reflects a considerable rej ection, the working of the principle was before us. Indeed, an old process of desilvering lead practised the principle by mechanically transferring the impoverished solid crystals in one direction and the enriched liquid in the other, through a series of crucibles which merely fluctuated, each over its own small range of temperature. The present embodiment illustrated schematically in Fig. 45, was developed by W. G. Pfann at the Bell Telephone Laboratories.86 in the purification of germanium the impurities, such as phosphorous, antimony and arsenic, are thought to be reduced to 1 atom in 10,000,000,000. No independent analytical means are available for such concentration, and the figures are derived by extrapolation from certain behaviours. The full effect of zone melting is manifest only when a series of many melted zones are passed along the bar in a refractory trough .

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(b)

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Fig. 45

Schematic illustration of zone refining process: (a) impure germanium ingot; (b) start of zone refining; (c) ultra pure germanium emerges (W. G. Pfann, Bell Telephone Laboratories).

The concentration of impurity at any distance along the zone melted bar, after successive passes, is shown in Fig. 46. To overcome the limitation of batch processes a new extension of the zone melting principle has been developed'? in which the raw, impure material is continuously

82

Hatfield Memorial Lectures VoL II

introduced at a midpoint and from which purified material is continuously withdrawn, as shown schematically in Fig. 47. The rejective influence is supplied by external moving heaters and since the principle of reflux is embodied in the method, this so called zone void method is really a counterpart, in crystallisation, of continuous fractional distillation. r- -

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Schematic diagram of continuous zone refining (W. G. Pfann, Bell Telephone Laboratories).

THERMODYNAMICS

IN METALS

It is generally recognised that a knowledge of the thermodynamic activity of the components of alloys and carbides, and also of slagsand mattes, is essential to an understanding of

Trends in Metallurgical Research. in the United States

83

the metallurgical reactions involving them. Existing .methods for measuring activities, such as gaseous equilibria, distribution or electromotive force studies, fail in many cases to supply the necessary data. The Knudsen cell is now being used successfully by Dr Law McCabe of Carnegie Institute of Technology to measure vapour pressures, from which activities can readily be calculated.s" The simplified, schematic diagram of the Knudsen cell in Fig. 48 shows that it is simply a cylinder with a small knife edged orifice in the top. The solid or liquid to be studied is placed inside the cell, heated in a vacuum to induce its volatilisation. From the amount of the component which escapes through the orifice in unit time, the pressure can be calculated from the Knudsen equation, which is derived in a straightforward manner from the Kinetic Theory of Gases. The amount of material effusing from the cell has usually been determined by the weight loss of the cell or by radioactive tracer techniques. The latter method allows pressures as low as 10-9 atmospheres to be measured with ease, tremendously extending the versatility of the Knudsen cell.

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Fig.48

Schematic diagram of Knudsen cell (L. C. McCabe, Carnegie Institute of Technology).

The thermodynamic activity of a component in a solution is equal to its pressure above the solution divided by the vapour pressure of the pure substance. Using this simple ratio, the activity of the following components has been determined: Cr in Fe-Cr; Ag in AuAg; Mn in Mn7C3; Mn in (Mn,Fe)7C3; Si02 in liquid CaO-Si02; Cu in liquid Cu2S; S in FeS (in equilibrium with Fe). Figure 49 shows the activity of chromium as a function of its concentration in iron-chromium alloys''? at 2192°F (1200°C). Another approach being used in obtaining quantitative thermodynamic data for metals and alloys involves the use of a high temperature galvanic cell. Cohen and co-workers?" have used the technique in studying the thermodynamic properties of gold-nickel, aluminium-zinc and aluminium-silver alloys. Activity curves at 1652°F (900°C) for gold-nickel alloys as determined by Seigle, Cohen, and Averbach are shown in Fig. 50, wherein large positive deviations from Raoult's law are apparent. Darken''! has derived

84

Hatfield Memorial Lectures VoL II 1·0

/

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Activity of Cr in Fe-Cr alloys at 1200°C (L. C. McCabe, Carnegie Institute of Technology) .

the following equation which connects the activity and the diffusion rates in a given system: dIna D = (N1D2 + N2D1) -al nNl

where D = interdiffusion coefficient, Nl and N2 = mole fractions of components 1 and 2, Dl and D2 = self-diffusion coefficients of components 1 and 2, a1 = thermodynamic activity of component 1.

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Fig.50

Activity curves at 900°C for Ni-Au alloys (L. L. Seigle, M. Cohen and B. L. Averbach, Massachusetts Institute of Technology).

Trends in Metallurgical Research in the United States

85

Inasmuch as the thermodynamic factor (dlna1)/(dlnN1) is now known for the goldnickel system as a function of composition (Fig. 50) it becomes possible to test the Darken equation through measurements of the interdiffusion coefficient and the self-diffusion coefficients* as a function of composition. This has now been done, with the satisfying agreement shown in Fig. 51.92 The 'calculated' curve was obtained from the right hand side of the Darken equation, whereas the 'observed' curve was determined experimentally by the Matano method.

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Calculated and observed interdiffusion coefficients in Ni-Au alloys at 900°C Reynolds and M. Cohen, Massachusetts Institute of Technology).

a. E.

It will be noted that there is a minimum in the diffusion rate corresponding to compositions at the top of the miscibility gap. In this region, the thermodynamic activity is insensitive to composition, and the above thermodynamic factor is quite low. Thus, the minimum diffusion rate results, not from low atomic mobility, but from a low driving force. Calorimetry provides still another method of obtaining thermodynamic data. Dr O. J. Kleppa of the Institute for the Study of Metals is engaged in a systematic calorimetric study of binary alloys among the group B metals using a highly refined calorimeter'< particularly suited for the study of heats of mixing and heats of solution in alloy systems. A cross-section of the calorimeter is shown in Fig. 52. Kleppa estimates that an accuracy of ± 1% in actual experimental observations can be achieved with his apparatus. *The self-diffusion coefficients were measured by means of the radioactive tracers Au 198 and Ni63.

86

Hatfield Memorial Lectures Vol. II

Fig. 52 Central part of calorimeter assembly: A is furnace core with main heaters; B is constant temperature jacket; C is heavy shield; D is calorimeter block; E is calibration inlet; F is removable crucible: G is charging and stirring device; H is inner plug; I is part of inner plug. (0.]. Kleppa, Institute for the Study of Metals).

PHYSICAL CHEMISTRY OF STEELMAKING Studies on the physical chemistry of iron- and steelmaking are widening the paths to more efficient and more intelligent operation of metallurgical processes. Of the investigators in this field, the outstanding and voluminous work of ProfessorJohn Chipman and his successive associates deserves special attention. Professor Chipman's contributions have brought not only enlightenment and understanding of the basic chemical and physical relationships involved in winning iron from its ores and the refinement of this iron into steel, but also have pointed the way toward further improvement in the technology of iron and steelmaking. It is reasonable to expect that metallurgists and physical chemists will continue for many years to refer to Professor Chipman's work on the free energy of iron oxides.?" the application of thermodynamics to the deoxidation of liquid steel.i" the equilibria and kinetics of slag-metal reactions involving oxygen,96 sulphur.P? phosphorus.?" and manganese."? More recently, his workl?" on hydrogen in liquid steel and slags promises to

Trends in Metallurgical Research in the United States 87 reveal the factors which control the behaviour of this most elusive element, and I am happy to present here a resume contained in a private communication from him. When a piece of steel containing hydrogen is cooled from the rolling temperature, the hydrogen solubility and rate of diffusion both decrease. At low rates of diffusion, the hydrogen is effectively trapped inside the steel in amounts which exceed its solubility in the solid metal. A large effective pressure can be developed which may be reflected in cracks or flakes, as suggested by Lukemeyer-Hasse and Schenck.tv! Zapffe and Sims102 proposed that hydrogen collects in the substructure as a means of accounting for the fact that hydrogen decreases ductility very markedly without having a corresponding effect on the hardness of the steel. Calculations showing enormous hydrogen pressures at room temperature require a slight modification. The higher pressures calculated are based on the ideal gas equation for hydrogen, which is not strictly valid at these extreme pressures. At room temperature and very high pressures, it is known that the actual pressure is substantially less than the fugacity or thermodynamic pressure which is derived from the calculations. Using the best estimate of the relation between fugacity and pressure for hydrogen at room temperature, it was shown by Carney, Chipman and Grant103 that a steel containing 5 ppm of hydrogen at room temperature could develop a hydrogen pressure of 12,000 atm. When an ordinary sample is taken from the steel bath, a substantial portion of its hydrogen content is lost. For this reason it is necessary that special methods be developed for securing samples of hydrogen for analysis. The sampling device which Chipman and associates have used successfully is a modified Taylor sampler illustrated in Fig. 53, which is taken from a paper by Epstein, Walsh and King. 104 In operation a well slagged spoon is used to remove a sample of metal from the furnace. The metal is killed in the spoon with about 0.2% aluminium, and the sample is sucked up into the copper mould where it solidifies very rapidly. The sampler is then supported on a tripod, and the sample is knocked out by means of the steel punch into a pail full of water. The time required for sampling and quenching the sample into cold water is less than 5 s. After about lOs agitation in the water quench, the sample is placed in liquid nitrogen and is maintained at this temperature until time for analysis. Experience in the development of this sampling device has pointed to one or two critical requirements which must be observed. The metal must be quenched very drastically from the liquid state into the gamma region to prevent the rapid diffusion which occurs in delta iron. It is then cooled to liquid nitrogen temperature as rapidly as possible, removed briefly to hammer off the ring of steel that solidified on the outside of the sampler, and returned to liquid nitrogen for storage. The determination of hydrogen in a steel sample by solution in liquid tin and analysis of the evolved gas has become a well established method. A recent addition to this method introduced by Shields, Chipman and Grant10S is the use of instrumental means for determining the amount and composition of the evolved gas. The gas is collected at a pressure of 1-2 mm Hg and is introduced into the thermal conductivity cell or 'hot wire gauge' shown in Fig. 54. The fine platinum filament is heated by a current whose strength is adjusted to give a predetermined resistance of the wire, that is, to a fixed wire

88

Hatfield Memorial

Lectures VoL II

Rubber gasket

Nipple

Fig. 53 Modified Taylor sampler for obtaining samples from liquid steel for hydrogen analysis (D. J. Carney, J. Chipman and N. J. Grant, Massachusetts Institute of Technology).

temperature. The current required to bring the wire up to this temperature depends upon the thermal conductivity of the gas, which is highly sensitive to its hydrogen content. Moreover, the two impurity gases which are present, carbon monoxide and nitrogen, have almost identical thermal conductivities so that it is unnecessary to know the relative proportions of these two gases. The hydrogen content of the sample is thus determined by the pressure, the reading on the burette, and the ammeter. It has become clear that most of the hydrogen which is found in steel came from water vapour rather than directly from gaseous hydrogen. The open hearth flame normally contains a very substantial proportion of water vapour, formed from the combustion of hydrocarbon fuels or resulting from steam used in atomising the fuel oil. Burnt lime normally contains considerable quantities of water unless it is very freshly burned or protected from the atmosphere between the time ofbuming and time of addition to the furnace. Water vapour is normally present also in the atmosphere, particularly in summer. Water vapour is able to convey not only hydrogen but also oxygen to liquid steel, and the chemical balance between the two requires that steels which are low in oxygen can absorb larger quantities of hydrogen. Equilibrium conditions in this process have been studied, and hence the limiting quantities of hydrogen which can be absorbed from water vapour are known and can be expressed in terms of the oxygen content of the liquid

Trends in Metallurgical Research in the United States 89

,

to bridQe to furnace

brass connectors tunQsten rods Qroundjoint sealed in wax

ice and water nickel wire

MERCURY CUT-OFF

fine platinum wire

I

wire loop

fine platinum . ..- connecting wire weight

THERMAL CONDUCTIVITY CELL

Fig. 54

Thermal conductivity cell for analysis of hydrogen (B. Shields, J. Chipman and N. Grant, Massachusetts Institute of Technology).

metal. The relation is shown in conditions of low relative humidity hydrogen if its oxygen content is susceptible to hydrogen absorption

J.

Fig. 55. From this figure it is evident that even under liquid steel is capable of absorbing a substantial quantity of low. Thus steel which has been deoxidised is particularly and must be guarded against the access of water vapour.

100 50

E

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z

tJ

10

0 ex: 0

>:t:

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Relation between water vapour pressure, oxygen and hydrogen in liquid metal Chipman, Massachusetts Institute of Technology).

O.

90

Hatfield Memorial Lectures VoL II CONCLUSION

Finally, another trend, one which I note here with special pleasure and gratitude, is the better acquaintance of British and American metallurgists, and the growing volume of exchange of research experiences, even in the early stages. Y ou and your colleagues in the United Kingdom have received us most warmly on our visits which, as it happens, have been somewhat more numerous than yours to us. I sincerely hope you feel that you are more than welcome in our laboratories and meetings, for indeed you are.

ACKNOWLEDGE~NTS This summary of the trends of some metallurgical researches in the many fields being investigated in America today could clearly be made only with the generous aid of many whose names appear in the references. To them I express my sincere thanks and nly appreciation also of their sympathetic interest in making this review timely and informative. Special thanks are due to my associate, Dr W. T. Lankford, jun.

REFERENCES 1. 2. 3. 4.

5. 6.

7.

8.

E. C. BAIN and E. S. DAVENPORT: Trans. An1. Inst. Mitt. Met. Eng., 1930, 90, 117. E. S. DAVENPORT: Trans. Ant. Soc. Met., 1939,27,837. G. M. ROBERTSON: Trans. An1. Inst. Min. Met. Eng., 1930,90, 150. (a) Atlas if Isothermal Transformation Diagrams: United States Steel Corporation, 1951, Pittsburgh. (b) Supplement to Atlas if Isothermal Transformation Diagrams, United States Steel Corporation, Pittsburgh 1953. B. L. AVERBACHand M. COHEN: Trans. Aln. Soc. Met., 1949,41,1024. (a) Fracturing if Metals, American Society for Metals, Cleveland, 1948. (b) Cold Working if Metals, ASM, Cleveland, 1949. (c) Thermodynamics in Physical Metallurgy, ASM, Cleveland, 1950. (d) Atom Movements, ASM, Cleveland, 1951. (e) Metal Inteifaces, ASM, Cleveland, 1952. (f) Modern Research Techniques in Physical Metallurgy, ASM, Cleveland, 1953. (g) Relation if Properties to Microstructure, ASM,· Cleveland, 1954. (h) Impurities and lmpefeaions, ASM, Cleveland, 1955. (i) Theory of Alloy Phases, (to be published). (a) H.]. GOLDSCHMIDT:]. Iron Steel Inst., 1948,160,345. (b) H.]. GOLDSCHMIDT:]. Iron Steel Inst., 1952,170, 189. (c) H.]. GOLDSCHMIDT:Symposium on High Temperature Steels and Alloys for Gas Turbines, Spec. Rep. No. 43, The Iron and Steel Institute, London, 1951,249-257. (a) H. H. LESTERand R. H. ABORN: Army Ordnance, 1925-1926, 6, 120, 200, 283, 364. (b) C. S. BARRETT: Structure if Metals, Chapter XIV: McGraw-Hill, New York, 1952.

Trends in Metallurgical Research in the United States

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9. J. T. NORTON: Private communication. 10. J. T. NORTON: Private communication. 11. J. T. NORTON: Private communication. 12. H. C. GATOS and AHMED AZZAM: Trans. Am. Inst. Min. Met. Eng., 1952, 194,407. 13. A. D. KURTZ, B. L. AVERBACHand M. COHEN: 'Self-Diffusion of Gold in Gold-Nickel Alloys' submitted to Am. Inst. Min. Met. Eng. 14. R. S. RUNDLE, C. G. SHULLand E. O. WOLLAN: Acta Crystall., 1952,5,22. 15. C. G. SHULLand S. SIEGEL:Phys. Rev., 1949,75, 1008. _ 16. C. G. SHULL,E. O. WOLLAN and W. C. KOEHLER: Phys. Rev., 1951,84,912. 17. Symposium on Fluorescent X-ray Spectrographic Analysis, Spec. Tech. Publ. No. 157: American Society for Testing Materials, Philadelphia, 1954. 18. L. S. BIRKS and E.]. BROOKS: Anal. Chem., 1955,27, (3),437. 19. C. S. BARRETT and O. R. TRAUTZ: Trans. Am. Inst. Min. Met. Eng., 1948,175,579. 20. C. S. BARRETT: Am. Mineralogist, 1948,33,749. 21. C. S. BARRETT: Private communication. 22. C. S. BARRETT: Private communication. 23. C. S. BARRETT: 'Metallographic Study of Sodium, Potassium, Rubidium, and Cesium after Cooling to 1.20 K.,' submitted for publication to J. Inst. Met. 24. J. C. FISHER,]. H. HOLLOMON and D. TURNBULL: Trans. Am. Inst. Min. Met. Eng., 1949, 185,691. 25. M. COHEN, E. S. MACHLIN and V. G. PARANJPE: 'Thermodynamics of the Martensite Transformation'in Thermodynamics and Physical Metallurgy, ASM, Cleveland, 1950. 26. S. A. KULIN and M. COHEN: Trans. Am. Inst. Min. Met. Eng~, 1950, 188,1139-1143. 27. (a) K. E. CECH and]. H. HOLLOMON: Trans. Am. Inst. Min. Met. Eng., 1953,197,685. (b) ]. C. FISHER: Acta Met., 1953, 1,32. 28. C. H. SHIH, B. L. AVERBACHand M. COHEN:]. Met., 1955,7,183. 29. E. S. MACHLIN and M. COHEN: Trans. Am. Inst. Min. Met. Eng., 1952,194,489. 30. ]. R. PATELand M. COHEN: Acta Met., 1953, 1,531. 31. A. H. GEISLER:Acta Met., 1953,1,260. 32. M. S. WECHSLER, D. S. LIEBERMANand T. A. READ: Trans. Am. Inst. Min. Met. Eng., 1953, 197, 1503-1515. 33. E. C. BAIN: Trans. Am. Inst. Min. Met. Eng. 1924,40,25. 34. See bibliographies for Refs 5,35,36,37, 38 and 39. 35. R. W. BALLUFFI,M. COHEN and B. L. AVERBACH: Trans. Am. Soc. Met., 1951,43,497. 36. K. H.]ACK:J. Iron Steel Inst., 1951, 169,26. 37. C. S. ROBERTS, B. L. AVERBACHand M. COHEN: Trans. Am. Soc. Met., 1953,45,576. 38. (a) B. S. LEMENT,B. L. AVERBACHand M. COHEN: Trans. Am. Soc. Met., 1954,46, 851. (b) B. S. LEMENT,B. L. AVERBACHand M. COHEN: Trans. Am. Soc. Met., 1955,47,291. 39. M. COHEN: Private communication (see also Ref. 38a). 40. M. COHEN: Private communication. 41. A. G. QUARRELL:Private communication. 42. E. C. BAIN and Z. JEFFRIES:Iron Age, 1923, 112, 805. 43. R. S. DEAN and B. SILKES:'Boron in Iron and Steel,' Information Circular 7363, Bureau of Mines, United States Dept. of the Interior, 1946.

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44. (a) R. A. GRANGE and T. M. GARVEY: Trans. A,ner. Soc. Met., 1946,37,136. (b) G. D. RAHRER and C. D. ARMSTRONG: Trans. Am. Soc. Met., 1948, 40, 1099. (c) R. A. GRANGE, W. B. SEENS,W. S. HOLT and T. M. GARVEY: Trans. Am. Soc. Met., 1950, 42, 75. (d) R. A. GRANGE and T. M. GARVEY:Steel, 1946,118,98. 45. M. E. NICHOLSON: Private communication. 46. M. E. NICHOLSON: Private communication. 47. R. M. FISHER: Symposium on Techniques for Electron Metallography, Spec. Tech. Publ. No. 155, ASTM, Philadelphia, 1953,49. 48. A. HULTGREN: Trans. Am. Soc. Met., 1947,39,915. 49. G. W. RATHENAU and G. BAAS:Physica, 1951,17,117. 50. (a) R. D. HEIDENREICH:]. Appl. Phys., 1955,26,757. (b) R. D. HEIDENREICH:]. Appl. Phys. 1955,26,879. 51. N. H. NACHTRIEB and G. S. HANDLER: Acta Met., 1954,2,797. 52. A. H. COTTRELL: Dislocations and Plastic Flow in Crystals, The Clarendon Press, Oxford, 1953. 53. F. C. FRANK: 'Advances in Physics', Phil. Mag., Supplement 1, 1952,1. 54. W. T. READ, jun.: Dislocations in Crystals, 1953, McGraw-Hill, New York, 1953. 55. W. SHOCKLEY:Trans. Am. Inst. Min. Met. Eng., 1952,194,829. 56. E. R. PARKER and]. WASHBURN: Modern Research Techniques in Physical Metallurgy, ASM, Cleveland, 1953, 186. 57. E. R. PARKER and]. WASHBURN: Impurities and lmpeijeaions, ASM, Cleveland, 1955, 145. 58. (a) W. G. PFANN and L. C. LOVELL:Private comminication. (b) F. L. VOGEL:Acta Met., 1955,3,245. 59. ]. GILMAN:Private communication. 60. (a) C. G. DUNN and W. R. HIBBARD:Acta Met., 1955,3,409. (b) C. G. DUNN and W. R. HIBBARD: Private communication. 61. W. I. POLLOCKand R. F. MEHL: Acta Met., 1955,3,213. 62. M. SIMNAD:Private communication. 63. K. G. COMPTON, A. MENDIZZA and S. M. ARNOLD: Corrosion, 1951,7,327. 64. S. ELOISEKOONCE and S. M. ARNOLD:]. Appl. Phys., 1954,25, 134. 65. R. M. FISHER, L. S. DARKEN and K. G. CARROLL: Acta Met., 1954,2,368. 66. S. ELOISEKOONCE and S. M. ARNOLD:]. Appl. Phys., 1953,24,365. 67. G. W. SEARS:Acta Met., 1955,3, (4), 367. 68. G. PFEFFERKORN:Z. Microscopie, 1955, 62, 109. 69. (a) C. HERRING: Bell Lab. Record, 1955, 33, 285. (b) C. HERRING and]. K. GALT: Phys. Rev., 1952,85,1060. 70. (a) G. W. SEARS,A. GATTI and R. L. FULLMAN:Acta Met., 1954,2,727. (b) G. W. SEARS:Gordon Conference,]uly, 1955. (c) A. W. COCHARDT and H. WIEDERSICH:Natunoissensthcften, 1955,42,342. (d) R. EISNER: Acta Met., 1955,3, (4),414. 71. G. W. SEARS:Gordon Conference,]uly, 1955. 72. (a) M. S. HUNTER and W. G. FRICKE,jun.: Proc. Am. Soc. Test. Mat., 1954,54,717.

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(b) M. S. HUNTER and W. G. FRICKE,jun.: 'Effect of Alloy Content on the Metallographic

73. 74.

75.

76.

77.

78. 79.

80.

81. 82. 83. 84. 85. 86. 87. 88. 89.

Changes Accompanying Fatigue' (To be published in Proc. Am. Soc. Test. Mat., 1955, 55). V. PnsPANEN:]. Appl. Phys., 1948,19,876. (a) M. E. MERCHANT:]. Appl. Phys., 1945,16,267. (b) M. E. MERCHANT:]. Appl. Phys., 1945,16,318. (c) M. E. MERCHANT and N. ZLATIN: Trans. Am. Soc. Met., 1949,41,647. (a) M. C. SHAW, N. H. COOK and I. FINNIE: Trans. Am. Soc. Mech. Eng., 1953,75,273. (b) M. C. SHAWand 1. FINNIE: Trans. Am. Soc. Mech. Eng., 1955, 77, 115. (c) I. FINNIE and M. C. SHAW: 'The Friction Process in Metal Cutting': Paper No. 54-Ai 08 presented at the 1954 Annual Meeting of the American Society of Mechanical Engineers. F. W. BOULGER, H. E. HARTNER, W. T. LANKFORDand T. M. GARVEY: 'Force Relationships in the Machining of Low-Carbon Steels of Different Sulfur Contents' (to be submitted for publication to the American Society of Mechanical Engineers). (a) B. T. CHAO, K.J. TRIGGER, and L. B. ZYLSTRA:Trans. Am. Soc. Mech. Eng., 1952,74, 1093. (b) B. T. CHAO and K. J. TRIGGER: 'Temperature Distribution at the Tool-Chip Interface in Metal Cutting', Paper No. 54-A-115 presented at the 1954 Annual Meeting of the American Society of Mechanical Engineers. T. S. ROBERTSON:]. Iron Steel Inst., 1953,175,361. (a) M. L. WILLIAMS:Symposium on Effect of Temperature on the Brittle Behavior of Metals with Particular Reference to Low Temperatures, Spec. Tech. Pub!. No. 158, ASTM, Philadelphia, 1954. (b) M. L. WILLIAMS:Investigations of Structural Failures in Welded Ships, National Bur. of Standards Report 4168, June 30, 1955. (a) P. P. PUZAK, M. E. SCHUSTERand W. S. PELLINI: Welding]., 1954,33, 481s. (b) P. P. PUZAK, M. E. SCHUSTERand W. S. PELLINI: Welding]., 1954,33, 433s. (c) P. P. PUZAK and W. S. PELLINI: Effect of Temperature on the Ductility of High Strength Structural Steels Loaded in the Presence of Sharp Cracks, Naval Research Laboratory Report 4545, June 22, 1955. (d) W. S. PELLINI: Symposium on Effect of Temperature on the Brittle Behavior of Metals with Particular Reference to Low Temperatures, Spec. Tech. Publ, No. 158, l\STM, Philadelphia, 1953,216. G. R. IRWIN: Fracturing of Metals, ASM, Cleveland, 1948. J. R. Low, jun.: Relation of Properties to Microstructure, ASM, Cleveland, 1954. J. R. Low, jun.: 'Dislocations and Brittle Fracture in Metals,' to be presented at International Union of Theoretical and Applied Mechanics meeting, Madrid, Sept., 1955. L. C. CHANG and M. GENSAMER:Acta Met., 1953, 1,483. M. GENSAMER:Private communication. W. G. PFANN: Trans. Am. Inst. Min. Met. Eng., 1952, 194,747. W. G. PFANN:]. Met., 1955,7,297. L. C. MCCABE:]. Met., 1955, vol. 7, Jan., p. 61. L. C. MCCABE: Private communication.

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90. L. L. SEIGLE,M. COHEN and B. L. AVERBACH: Tra1ts. Ant. Inst. Min. Met. Eng., 1952, 194, 1320. 91. L. S. DARKEN: Trans. Am. Inst. Min. Met. Eng., 1948, 175, 184. 92. J. E. REYNOLDSand M. COHEN: Private communication. 93. O. J. KLEPPA:}. Phys. Chem., 1955, 59, 175. 94. J. CHIPMAN: Indust. Eng. Chem., 1933, March, 319. 95. J. CHIPMAN: Trans. Am. Soc. Met., 1934,22,385. 96. (a) J. CHIPMAN: Am. Inst. Min. Met. Eng.} Open Hearth Proc., 1942, Sept., 695. (b) S. MARSHALLand]. CHIPMAN: Trans. Ant. Soc. Met., 1942,30,695. 97. N.]. GRANT and]. CHIPMAN: Trans. Am. Inst. Min. Met. Eng., 1946, 164, 134. 98. T. B. WINKLER and]. CHIPMAN: Trans. Am. Inst. Min. Met. Eng., 1946, 164,111. 99. ]. CHIPMAN,]. B. GERO, and T. B. WINKLER: Trans. Am. Inst. Min. Met. Eng., 1950, 188, 341. 100. ]. CHIPMAN: 'Hydrogen in Steelmaking,' presented at a metting ofJernkontoret in Stockholm on May 21, 1955 (To be published inJernkontorets Annaler). 101. L. LUCKMEYER-HASSEand H. SCHENCK:Arch. Eisenhiittenwes., 1932. 6,209. 102. C. A. ZAPFFEand C. E. SIMS: Trans. Am. Inst. Min. Met. Eng., 1941,145,225. 103. D.]. CARNEY,]. CHIPMAN, and N.]. GRANT: Am. Inst. Min. Met. Eng., Electric Furnace Steel Proceedings, 1948, 6, 34. 104. H. EpSTEIN,]. H. WALSH, and T. B. KING: Reg. Tech. Meet., Ant. Iron Steel 1I1st., 1954. 105. B. SHIELDS,]. CHIPMAN, and N.]. GRANT: Trans. Am. Inst. Min. Met. Eng., 1953, 197, 180. 106. (a) First Progress Report, Subcommittee XI of ASTM Committee E-4 on Metallography, Proc. Am. Soc. Test. Mat., 1950, 50, 444. (b) Second Progress Report, Subcommittee XI of ASTM Committee E-4 on Metallography, Proc. Am. Soc. Test. Mat., 1952, 52, 540. (c) Third Progress Report, Subcommittee Xi of ASTM Committee E-4 on Metallography, Proc. Am. Soc. Test. Mat., 1953,53,503.

THE

TENTH

HATFIELD

MEMORIAL

LECTURE

The Mechanism of Formation of Banded Structures Paul G. Bastien, D.Sc. At the time the lecture was given Professor Bastien was Scientific Director at the Etablissements Schneider and consultant to the Commissariat a l'Energie Atomique. The lecture was presented at the Firth Hall, Sheffield University, on 28 November 1957.

When I received the invitation to give the 10th Hatfield Memorial Lecture, I felt both greatly honoured and at the same time rather apprehensive at seeing my name added to the list of eminent metallurgists who have been called upon to honour the memory of the late Dr Hatfield by delivering this Lecture. Some years before the second world war I had the opportunity to meet Dr Hatfield on several occasions, and I was very greatly impressed by the breadth of his knowledge, his depth of thought and his method of expression, which was always extremely clear, often even jocular, and always to the point. The choice of a subject for this Lecture, which should be connected with Dr Hatfield's work, is a relatively easy task, in view of the range and variety of his activities. For many years he served as the distinguished Chairman 'of the Heterogeneity of Steel Ingots Committee of The Iron and Steel Institute, and for that reason, and also because the 10th Hatfield Memorial Lecture is being given in Sheffield, a town well known in the world of metallurgy, it seemed appropriate to discuss on this occasion the mechanism of formation of the banded structure in steel, a subject in which I have taken a great interest over many years.

Manufacturing processes which involve plastic deformation without the removal of metal have an important effect on the structure of a metal or alloy, not only on the macro- and microscopic scale but also on the crystalline structure; furthermore, this effect differs greatly, depending on whether the deformation takes place at high or low temperature." Several phenomena are known to occur during the non-isothermal solidification of an alloy over the solidification range. In the case of steel, they cause the composition to vary in the different parts of the material and give rise to the following effects: Large scale segregation, variations in concentration between the liquid and solid phases during solidification which are found over the whole of the ingot. This has formed the subject of a large number of investigations, among them the outstanding studies of the

Heterogeneity of Steel Ingots Committee of The Iron and Steel Institute.f

95

96

Hatfield Memorial Lectures VoL II Small scale or dendritic segregation, which affects the crystals of primary crystallisation and

implies the existence of solid solutions in which the concentration varies with temperature. In a worked steel we have to consider the following structures resulting from the plastic deformation: • A macroscopic laminated structure caused by dendritic segregation • A microscopic laminated structure, which is particularly marked if forging takes place between the A3 and Al points or below the Al point • A crystalline laminated structure, which results from the actual mechanism of the plastic deformation, mainly slipping. One of these structures will predominate and control the properties of the final product, depending on the type of steel and its production, the method of casting, the weight and shape of the ingot, the actual conditions under which the plastic deformation takes place, in particular the range of temperatures and the amount of deformation, and finally the intermediate reheats and the final heat treatment.

THE BANDED STRUCTURE: DESCRIPTION AND CONDITIONS FOR ITS APPEARANCE During the solidification of steel, the most impure liquid with the lowest melting point, which contains in particular sulphur inclusions in the form of suspensions, becomes concentrated and trapped in the spaces within the dendritic network. It has been suggested, first by Stead-' and later by Howe," Sauveur'' and Whiteley," that during the cooling through the transformation range the phosphorus concentrated in the interdendritic spaces displaces the carbon from these areas towards the arms of the primary crystals. According to this view, the process results in the presence of dendrites whose arms are rich in carbon and poor in phosphorus, while the interdendritic spaces have a low carbon and a high phosphorus content. This effect, however, appears only in steels with a relatively high phosphorus content and for suitable rates of cooling through the transformation region. Several years ago we had the opportunity to examine a number of300 mm (12 in.) diameter forged cylinders after quenching and annealing, made from Ni-Cr steels produced in 6 ton ingots in acid, basic and duplex electric furnaces.? These steels had a low sulphur and phosphorus content (S 0.011-0.031%; P 0.015-0.021%) and a suitable cooling rate (water quench after a homogenisation reheat). It was found that because of the low phosphorous content there was no displacement of carbon towards the dendrite axes, the carbon being concentrated in the original interdendritic spaces in accordance with the usual effect of small scale segregation. 8 During solidification, the inclusions themselves are also subject to large and small scale segregation. The silicate inclusions are concentrated in the arms of the dendrites, giving a

The Mechanism of Formation of Banded Structures

97

certain amount of support to Benedick's theory? that inclusions rich in oxygen act as nuclei for crystallisation from the liquid during the solidification of the ingot. The sulphur inclusions, on the other hand, are located in the interdendritic spaces in accordance with the usual mechanism of small scale segregation. This small scale segregation would disappear if the inclusions could be absorbed by means of a long heat treatment at the high temperatures of the gamma region, thus making the material entirely homogeneous chemically. This is never the case in practice, mainly because of the low diffusion rates of phosphorus and most other dissolved elements (except carbon), and also because of the very low solubility of non-metallic inclusions in the solid state in iron. This segregation therefore persists, although to a more or less reduced extent, even after the high temperature homogenisation heat treatments used for large forgings. To simplify matters, let us consider an ingot hot worked by forging or rolling, after complete solidification but without return to ambient temperature. The plastic deformation has the effect of modifying the dendrites by aligning them and deforming them mainly in the direction of the metal's greatest elongation; it follows that the inclusions must tend to form layers or laminae. This results in a fibrous or laminated structure which, in the case of rolling or forging, can readily be demonstrated by the use of metallo grap hie reagents, such as cupric chloride or iodine (Fig. 1).

Fig. 1 Laminated structure of forging. Etched in Stead's reagent (xl).

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During cooling, the separation of the proeutectoid constituents depends on the local chemical composition and is therefore initiated by the dendritic segregations which have been arranged in an orderly manner during hot working. When there is a pronounced phosphorus segregation, a micrograph shows a banded structure in the deformed material (Fig. 2); the white bands correspond to interdendritic regions which, as shown, are in general rich in phosphorus and poor in carbon and have a relatively high inclusive density, particularly manganese sulphide. The presence of manganese sulphide inclusions and the phosphorus segregation appear to be related phenomena.

Fig. 2

Macroscopic banded structure. Etched in Stead's reagent (x l 00)

A high temperature reheat followed by rapid cooling reduces the intensity of the banded structure without suppressing it entirely. If cooling takes place slowly, the banded structure reappears and can be even more pronounced than before.P- This is because the phosphorus segregation is only reduced very slowly at high temperature and, if we accept Stead's theory for the moment, it continues to displace carbon from the bands with a high phosphorus content, an effect which becomes more pronounced as the rate of cooling through the Ar3-Arl transformation range is reduced. For that reason the effectiveness of a homogenisation treatment diminishes with increasing size of the forging, and especially near its centre. The segregation of ferrite in the case of hypoeutectoid steels (or of cementite in hypereutectoid steels) thus results from an excessively slow cooling through the Ar3-Arl range in the presence of a pronounced phosphorus segregation. The accumulation of inclusions in the interdendritic regions can perhaps also contribute towards the effect to some extent, with the interfaces between the inclusions and the austenite acting as

starting points for the growth of proeutectoid crystals.

The Mechanism of Formation of Banded Structures

99

During the investigation into the Ni-Cr steels with low phosphorus content, in which there was no carbon displacement towards the dendrite axes, we found that the banded structure can also appear in the course of heat treatment. The cooling rate of the forging during quenching decreases continuously from the skin towards the centre; whenever it has a suitable value at the instant when the transformations take place, the small scale segregation produces structural heterogeneity by the formation of bands which differ micrographically and which extend along the fibre of the forging (Figs 3 and 4). Structures characteristic of annealing, quenching and tempering, as well as intermediate structures, can occur together in the same microscopic 'field (Fig. 5). This simple result is very important and suggests already that phosphorus is not the only cause of banded structures but that all the elements which contribute to the production and growth of small scale segregation must influence this phenomenon. In the case of the Ni-Cr steels, the sulphur inclusions are always found in those regions which are the most heavily- quenched, because they are richest in carbon, and, in the absence of carbon displacement by phosphorus, correspond to the original interdendritic spaces during solidification (Fig. 4)."The silicates, on the other hand, are located in the bands of the annealed structure, within the arms of the original dendrites, and more particularly in the ferrite areas (Fig. 3). To summarise, it is evident that a macroscopic laminated structure is formed through the collapse of a heterogeneous dendritic structure even when forging entirely above the

Fig. 3

Banded structure with no carbon displacement; silicates between bands of ferrite (x l00).

100

Fig. 4

Hatfield Memorial Lectures VoL II

Banded structure with no carbon displacement; sulphur in carbon rich areas, characterised by bands of pearlite-ferrite aggregates (xl00).

Fig. 5

Pearlite and bainite in same macroscopic field (xt 00).

The Mechanism of Pormation of Banded Structures

101

A3 point. During subsequent heat treatment this gives rise to a banded structure consisting of layers of similar microscopic constituents with equiaxed crystals, which can readily be detected under the metallurgical microscope. The shape and mechanism of formation of these structures varies, depending on the chemical composition, the impurities, and, at least so it appears at first, on the phosphorus content and the cooling rate at the instant of the transformations in the area concerned. Let us now consider a forging worked below the Ar3 temperature. If we take the case of a hypoeutectoid steel, a deposit of proeutectoid ferrite is formed from the austenite as the temperature falls, and the plastic deformation contributes to the formation of laminre of these two constituents; during the passage through the Ar 1 point the austenite is converted to pearlite. If there is no recrystallisation of one of the constituents, pearlite, or of both pearlite and ferrite, the microstructure remains permanently deformed and shows lamina; or bands of constituents of the same type which extend in the direction of drawing and are in some cases markedly non-equiaxed (Fig. 6). In this way we obtain a different type of banded structure; it is only noted in passing and will not be dealt with in detail since it does not form the subject of this lecture.

Fig. 6

Cold worked O.S%C steel (after Krivobok).

The banded structure proper which is found in worked materials under certain conditions has been investigated especially for annealed hypoeutectoid carbon steels.l? It is particularly frequent in steels containing between 0.10 and O.35%C and appears only rarely in steels with more than 0.6%C. It is found more often in fine grained steels and it is favoured by austenitising slightly above the AC3 point or, better still, between AC3 and Ac1, while a high temperature austenitising treatment is sufficient to suppress it in certain cases. Since the banded structure disappears when the cooling rate exceeds a certain value, it has been proposed 11 to define the susceptibility of a particular steel to this type of structure by this 'critical cooling rate for disappearance' which varies from one steel to

another.

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Hatfield Memorial Lectures VoL II

The disappearance of the banded structure is only temporary, as we have seen, since a further heat treatment followed by slower cooling causes it to reappear. Only a high temperature treatment over an extended period can permanently reduce to some degree the tendency of a steel to form a banded structure by reducing the dendritic segregation through diffusion. Numerous investigations on this subject12-15 have shown that this result is very difficult to achieve and that it is necessary in some cases to raise the temperature to 1300°C which, of course, raises other difficulties. Finally, certain workers!" have noted a relationship between the width of the bands and the temperature necessary to suppress them (Fig. 7).

u o •• 12 OOt---+_--J---+---+---O--::bo~-+---t---r--i UJ

a::

:::)

~

ffi

c, L

1I00~---+-----I---7"'---t-'-----r---

...-

w

-+-4----io •

Banded structure eliminated ~ almost unchanqed II

II

0·10

0·20

0'30

0·40

THICKNESS OF BANDS, mm

Fig. 7

Relationship between thickness of bands and reheating temperature (after Thompson and Willows10).

DENDRITIC SEGREGATION IN STEEL: INVESTIGATIONS The considerations outlined previously indicate the part played by dendritic segregation in the formation of the banded structure and clearly show the great importance of an experimental investigation. Conventional methods, using Baumann prints or cupric or iodine reagents, provide only a qualitative picture. They make it possible to demonstrate the presence of segregation without indicating the nature of the elements involved or the extent of the phenomenon. Although chemical analyses can readily be used, they involve the development of methods for removing extremely small samples which have to be accurately located in relation to the dendrites. Finniston and Fearnehough 16 carried out a number of microchemical analyses, using absorptiometric methods, on samples removed by means of a 1;16 in. (1.6 nun) diameter drill. During an investigation at Le Creusot into the heterogeneity of a forged ingot, in collaboration with Cattier, Dubois and Bleton," we have been able to reveal the different segregated bands by a deep selective acid etch which produced a

The Mechanism of Formation of Banded Structures

103

series of ridges and hollows on the surface of the worked metal. By careful machining it was possible to obtain small quantities of material from the ridges which in this case corresponded to areas richest in the segregated elements. The operation can be repeated ten or fifteen times and produces a sufficiently large sample for the determination of several elements by conventional methods. The technique of spectrographic point analysis, as developed by Convey and Oldfield.l? was also used by Finniston and Fearnehough 16 for the investigation of small scale segregation; several metallic elements could thus be determined in a very small sample of the order of 5 X 10-6 g. A number of novel methods which are likely to find extensive use in the study of segregations are: • Microradiography • Quantitative autoradiography • X-ray fluorescence point spectrography. The last of these is of particular importance. Microradiography This method, which represents the use of X-ray radiography on the microscopic scale, employs a characteristic radiation Ka, chosen so that the various constituents to be investigated differ widely in their absorption coefficients for this radiation. The method has been employed mainly by Trillat.P' Homes and Gouzou.!? and Wolfe and Robinson-"; its first use for the study of segregations in steel appears to have been by Betteridge and Sharpe.e! Quantitative Autoradiography The element whose segregation is to be studied must previously have been activated in a neutron beam from a nuclear reactor. A suitable sensitive emulsion is then applied with certain precautions to the polished surface of the steel sample and the radiation emitted by the labelled element produces a so called autoradiographic image which indicates its distribution in the surface layer. It is possible to obtain some quantitative information on the distribution of the radioactive element by using a microphotometer and comparing the photographic density at various points of the picture of the sample with that taken from a standard specimen in which the labelled element has a known activity. Using this method Kohn+' investigated at IRSID the dendritic segregation of phosphorus and arsenic in steel, as well as its persistence during thermal cycles, and developed a special microphotometer for this purpose. X-ray Fluorescence Point Spectrography This method was developed by Castaing+' at the Office National d'Etudes et de Recherches Aeronautiques (ONERA). It consists ofprojecting a very fine beam of electrons

104

Hatfield Memorial Lectures VoL II

on to the sample and analysing the radiation emitted by the metal after the impact of the electrons, using an X-ray spectrometer. For any given element the intensity of a characteristic radiation emitted by the alloy can then be compared with the intensity of the same radiation emitted by the pure element, so that it is possible to determine the concentration of this element in the volume irradiated; in Castaing's electronic microanalyser this volume corresponds to a cylinder 2 urn in diameter and 1-2 J.lm high. All elements whose atomic number is greater than 17 can thus be determined with an accuracy of 1% for concentrations greater than 0.5%. Equipment of this type is being operated at IRSID and is used for investigations into segregates.P" Some of the results obtained in the course of an unpublished investigation into dendritic segregation in a large forging ingot, carried out jointly by IRSID and the laboratory of Le Creusot, are given below. It is not the intention here to cover the entire literature on the tendency of the various elements in steel to produce dendritic segregation. This would have to include experimental results which cannot readily be compared because the information published depends on

• The steel: chemical composition and impurity content, method of manufacture,

casting temperature and weight of ingot which determine the size of the dendrites in the primary crystallisation • Its treatment: working and reheating cycles which influence diffusion • The method of defining segregation: volume of sample removed, method of analysis • The method of defining intensity oj segregation: certain authors compare the proportions in the segregated regions with those in the areas containing the purest dendrites, others with the mean composition of the steel. The discussion is limited to a qualitative and necessarily very general comparison between the tendencies of the main elements to produce segregation, and an indication of some of the results obtained with Castaing's electronic probe. Among the non-metals, carbon,7,8,16,25,26 phosphorus7,8,16,22,26,27 and sulphur7,8,16,26 can be considered to have a marked ability to produce segregation, but the diffusion of one can be greatly affected by the presence of the others, as well as by other metal constituents; this is particularly the case with carbon in the presence of phosphorus. The existence of arsenic segregation can be demonstrated directly by autoradiography,28 but the intensity of this segregation produced by different methods has not yet been clearly established; 7,8,28 in addition, there is no agreement on the possibility of homogenising this type of segregation by a high temperature heat trcatment.Fv? Among the metals, manganese16,19,20,24,27,30 produces extensive segregation on the dendritic scale; copper-'! and tin,31 molybdenum.I" tungsten.s" vanadium-" and probably cobalt also exhibit appreciable segregation. Nickel? ,8, 16,24produces a fairly small dendritic segregation which, however, is clearly visible. Chromium 16,34 and silicon 16 generally produce only slight dendritic segregation.

The Mechanism of Formation of Banded Structures

105

In general it is true to say that the impurities and the alloying elements in steel are as a whole more or less subject to dendritic segregation, to an extent which depends on the nature of the element and on its proportion in the steel, other conditions being equal. Carbon occupies a special place owing to its high diffusion coefficient which makes it easy to achieve its homogenisation within the austenite by heating above the A3 point, provided that it has not become fixed in the interdendritic spaces in the form of carbide by elements such as chromium, tungsten, vanadium or molybdenum, and thus rendered rather difficult to redissolve. When the steel is colled slowly through the Ar3-Ar1 range, this homogeneity of the carbon does not persist, the segregated elements locally modifying the transformation temperatures, or, more generally, the type of decomposition of the austenite. We have tried to determine quantitatively the amount of dendritic segregation in a 100 ton ingot of a Ni-Cr-Mo steel of the following composition: 0.24 C - 0.30 Si - 0.024 S - 0.015 P - 0.53 Mn - 2.5 Ni - 0.65 Cr - 0.27 Mo - 0.30 Sn. A core of 17Y2 in. (435 mm) diameter was taken from the axis of a rough forging of diameter 7 ft 4 in. (1850 mm) produced from this ingot, and J. Philibert of IRSID investigated the segregation as'shown by the very characteristic local heterogeneity of the microscopic structure, using Castaing's electronic probe. The results given here are those obtained for two areas only. Area 1 (bottom of ingot, suiface of core). Microstructure: alternate light (ferrite) and dark (pearlite + ferrite) bands. The following results were obtained:

First band

Microstructure

Cr,%

Ferrite

0.66-0.73

{

Second band Third band Fourth band

{

+ Pearlite Ferrite Ferrite

Ni,%

}

1.98-2.12

0.73-0.80 0.58-0.63 0.64-0.68

1.65-1.80 1.90-2.00

0.66-0.75 0.58-0.64

2.00-2.20 1.65-1.80

+ Pearlite Ferrite

When these results are expressed as proportions, they indicate a segregation of 30% of both chromium and nickel; the difference in the nickel and chromium concentrations in the ferrites within the ferrite and the ferrite-pearlite bands is particularly interesting. Similar results were obtained from measurements of the concentration of manganese and molybdenum; for example, the concentration of molybdenum ranges between 0.15 and 0.35%, with the following probable mean values: 0.3% pearlite, 0.2% ferrite.

Area 2 (V segregates, 53% by weight from bottom of ingot). Microstructure

consisting of three constituents arranged in two types of zone: a bainite zone and a zone containing ferrite and pearlite. The following results were obtained (Fig. 8):

106

Hatfield Memorial Lectures VoL II Microstructure Bainite Pearlite

Cr,%

Ni,%

0.80-0.85 0.65-0.70

2.80

+

}

Ferrite

2.15-2.25

0.55-0.65

I~------~--------~------~------~ CO

~ o·q ~

0-8

~

0'7~ O·b

_ ~~

s5 0·5t-------~-------..J.--------'------____i c

Bainite

o

Pearlite

x

Ferrite

Lengths

About 100;;.

Fig. 8

Results of point analysis made with the Castaing microanalyser.

The more heavily quenched structure thus appears to be richer in special elements, with the relative segregation of nickel and chromium being about 35%. The complete results, only a small part of which have been reproduced here, indicate that dendritic segregation is usually of the order of 30% for nickel, chromium, molybdenum and manganese. In the large ingot used in this investigation they are thus of the same order as the large scale segregations for the ingot as a whole, as determined by the chemical analysis of drill samples. Despite the fact that Castaing's probe analyses only a very small volume of material (a few cubic micrometres), there has been no evidence of any pronounced maxima in dendritic segregation which differ appreciably from the average values. The most heavily quenched constituent is always found to be richest in alloying elements. Microsegregation always takes place on a very small scale, approximately 10 J..lmwhen a distinction has to be made between ferrite and pearlite in a pearlite-ferrite aggregate; this distance increases to 50 Jlm and more for the case of an overall differentiation between bainite, pearlite and ferrite.

THE ACTUAL MECHANISM OF FORMATION STRUCTURE

OF THE BANDED

The mechanism of formation of banded structures caused by carbon segregation in the presence of phosphorus will be dealt with first, since this appears to play an essential part in the case of hypoeutectoid carbon and low alloy steels.

The Mechanism of Formation of Banded Structures

107

The mechanism of the carbon migration may be described as follows. After solidification there are more carbon and phosphorus atoms in the interdendritic spaces than near the axes of the dendrites. It has been shown in pure iron-phosphorus alloys and in steels having increasing proportions of this element that the presence of phosphorus raises the temperature of the A3 point at which austenite is transformed into ferrite (this effect is not shown by sulphur, for example). During a slow cooling, which allows the austenite to transform in the upper range, a steel with a relatively high phosphorus content will begin to become ferritic at a higher temperature than a steel containing less phosphorus. The same considerations apply to two separate areas of the same steel which differ in their phosphorus content as a result of dendritic segregation; ferrite will begin to fonn where phosphorus content is higher, while carbon becomes concentrated in those regions which are still austenitic. The mechanism described above raises two important questions:

(a) JiVhy does the carbon diffuse away from the areas rich in phosphorus, and not the phosphorus? The reason is that the activation energy of carbon for diffusion in austenite is smaller than that of phosphorus (43,000 and 32,000 cal/mol for phosphorus and carbon respectively), so that the carbon atoms diffuse more readily.

(b) JiVhy does the carbon migration result in an inversion of the heterogeneity and not in a uniform distribution of this element? It has already been shown, especially by Rey,32 that diffusion is controlled not by the concentration gradient but by an activity gradient. Numerous tests have indicated that the phosphorus reduces the solubility of carbon in alpha iron, thereby increasing the thermodynamic activity of the former. Atoms thus diffuse from regions of high activity to regions of low activity, even if this leads to an increase in concentration (uphill diffusion). Since the carbon migration is connected with the transformation from austenite to ferrite at A3 in a heterogeneous metal, it is not surprising that the migration is not observed when the upper transformation disappears owing to sufficiently rapid cooling. It is possible to extend these considerations still further and to apply them to practical cases. According to Grossmann.V phosphorus in solid solution produces a relatively large increase in the hardenability of the steeL But if there is both a rise in the A3 transformation temperature and an increase in hardenability, the curves representing the start of the austenite transformation in steels of different phosphorus contents must intersect (Fig. 9). Curve 1 shows, as a function of temperature, the time after which austenite begins to decompose in a very low phosphorus steel; curve 2 applies similarly to a steel with a higher phosphorus content. These two curves intersect at x owing to the effect described previously. It is obvious that the two curves could apply not only to two different steels but also to two different regions of the same steel: one corresponds to the dendrite axes (1) while the other is affected by segregation and is richer in phosphorus (2). If cooling takes place slowly (curve (a)), as in the case of a reheat treatment, the cooling curve would intersect curve 2 at M, and curve 1 at N at a lower temperature. Ferrite is first deposited in regions with a higher phosphorus content, resulting in concentration of carbon in the remaining austenite.

108

Hatfield Memorial Lectures VoL II Slow /cooling

......~ -:. - - ~ - .sc:>

.>,

Curve of deposition of proeutectold ferrite

<,

Quicker cooling

"'}. N

X

'\'. Low phosphorus

content

(b>

\

\ \ I I I

,

I I I

,I : (0)

LOG TIME

Fig. 9

Diagram showing conditions for eventual migration

of carbon.

With a more rapid rate of cooling, the corresponding cooling curve (b) intersects the transformation curves (1) and (2) at N' and M'. The austenite transformation now begins in regions with a low phosphorus content and tends to produce a concentration of the carbon in areas rich in phosphorus. The usual migration of the carbon away from phosphorus rich areas will therefore not take place in this case. These considerations lead to the conclusion that, from a certain cooling rate onwards, the regions with the highest phosphorus content can also be the last to begin the allotropic transformation, thus resulting in a concentration of the carbon atoms in regions richest in impurities. Provided that the composition of the steel is suitable, it should therefore be possible by means of a heat treatment, followed by cooling at a suitable rate, to induce the carbon atoms to return to the original interdendritic spaces from which they have diffused during the preceding sufficiently slow reheat treatment. These various cases were in fact found to be present during the comparative investigation into the banded structure of a Ni+Cr forging produced by acid, basic or duplex methods.o" as described above. In order to provide an even more conclusive confirmation of the mechanism described above, it would be useful to establish experimentally the displacement of the transformation curves under the action of the phosphorus. This was in fact done in the Research Laboratories of the Le Creusot works in the course of a general study-" of the banded structure of a steel of composition C, 0.28%; Ni, 1.56%; Cr, 0.52%; AI, 0.008%, to which phosphorus had been added to the extent of 0.025% (I), 0.056% (II) or 0.128% (III). The effect of the phosphorus on the temperature of the AC1 and Ac , points is shown in Fig. 10, which clearly indicates that the A3 point is raised under the influence of the phosphorus.

The Mechanism of Formation of Banded Structures

/

sse

soo

.: +.

v

W 850 0:: ::>

< ffi

.-

0-

~

109

-: 800 +

7001~--~O~.1~0----~O~·2~0----O~·~30~----~ PHOSPHORUS

J

Fig. 10

%

Influence of phosphorus on the position of the AC1 and AC3 points.

Transformations under non-isothermal conditions were studied mainly during steady cooling and by comparing the steels containing 0.025% (I) and 0.128% (III) of phosphorus; this could be done very accurately by means of a Chevenard differential dilatometer, using instead of the standard bar a reference bar of steel I, and placing the sample of the steel to be investigated (steel III) in the usual position.P" This arrangement gave very interesting results which have already been discussed elsewhere; it will be sufficient to note here that on cooling, the transformation in the test sample starts before that in the reference sample if the record curves upwards, and later in the test bar if the graph curves downwards (Fig. 11). Similarly, the transformation is completed in the test bar earlier or later than in the reference bar if the curve returns to the horizontal from above or from below. Conventional dilatometer tests (samples I and III compared with a reference bar), as well as the more sensitive tests described above (sample I compared with sample III), have given the results shown in Fig. 11, which indicate the effect of phosphorus on the gamma ---7 alpha transformation during cooling. The mean cooling rates between 850°C and the completion of the transformation are shown in Table 1. Table 1

per hour per minute per second

Mean cooling rates,

a

b

c

d

e

50

100

150

300

450 7.5

13

°c 9

h

50

90

225 3.75

12

110

Hatfield Memorial Lectures VoL II

Amphfication 180

1

215

bOO

700 ~ REGION

Fig. 11 Differential dilatometry of phosphorus steels. Reference specimen (1), O.025%P; comparison specimen (III), 0.128%P . It is evident that the phosphorus raises the temperatures of both the beginning and the end of the transformation for the lower cooling rates a-g, whilst it lowers them for the very high cooling rate j. For an intermediate speed of cooling (I), the transformation starts first and finishes last with a steel having a high phosphorus content. These experimental results thus confirm the basic principle of Fig. 9, namely the existence of a point of intersection x between the upper branches of the transformation curves of two steels with different phosphorus contents, and also possible mechanisms for separation of the proeutectoid ferrite described above. A useful addition to this information is provided by a microscopic study of isothermally transformed steels, which indicates that the phosphorus does not appear to have any effect on the position of the M, point (255°C), while an increase in the proportion of this element displaces the bainitic region towards the right and downwards. The beginning of the bainitic transformation is delayed slightly, and the actual transformation is slowed down considerably in the case of high phosphorus contents, mainly near 500°C. These results lead to a better explanation of the various structures found in the bands. In the case of the two steels discussed above (I and III), the intersection x of the curves for the start of the austenite ---7 ferrite transformation corresponds to a cooling rate of about 4°C per second. This eventually leads after a very limited ferrite separation to a structure

The Mechanism

of Formation

of Banded Structures

111

which is almost entirely bainitic and very unsuitable for the occurrence of banded structures. In other words, the increase in hardenability due to phosphorus in the steels investigated is too small and therefore fails to produce the conditions where the ferrite bands coincide with the original dendrite axes, the opposite of the usual result.

Effect of Dendritic Segregation of Elements other than Phosphorus The conclusions reached above suggest that the various elements capable of producing dendritic segregation in steel can either reinforce or counteract the effect caused by phosphorus, as has already been proposed by various authors. 12,13,36 According to Chipman.V the segregation tendency of an element L is proportional to the 'segregation coefficient' 1 - k, where

k (XOIo) sand

=

(XOIo)s/(XOIo)[

(XOIo) 1 being the proportions

of the element in the solid and in the liquid, according to the Fe-L equilibrium diagram. If Raoult' s law applies both to the liquid and to the solid, and if the element is present in a low concentration, Chipman shows that

l-k=

MXdT 1000 X%

where M is atomic weight of element dissolved in iron; )(010 is proportion of element dissolved in iron; ~ T is lowering of melting point due to an )(010 concentration of element. Values for 1 - k calculated by Chipman are shown, together with those for arsenic and tin:35 0.87C - 0.87P - 0.98S - 0.9802 -,..., lAs -,..., 0.7Sn - 0.34Si - 0.16Mn - O.44Cu0.20Ni - 0.05Cr - 0.20Mo. The impurities in steel (phosphorus, arsenic and tin) all suffer a very pronounced segregation and also have a similar effect on the temperature of the A3 point, which they increase appreciably. In view of their slow diffusion in iron, it might be thought that they play a similar and essential part in the formation of the banded structure, particularly in the case of low alloy or carbon steels. In the course of an investigation at present in progress at the laboratories of the Le Creusot works, we have studied the effect of various tin contents (0.035%, 0.125% and 0.260%) on a Ni-Cr steel of the same type as that used for the phosphorus investigation. The results obtained so far confirm that tin has a similar effect to phosphorus but to a lesser degree. It produces an increase in the temperature of the A3 point (Fig. 12), but smaller than that found in the previous investigations; as a result the transformations from the austenitic state during cooling start earlier in a steel with a higher tin content (Fig. 13). As early as 1918, Ie Chatelier and Bogitch-" suggested that the oxygen dissolved in steel plays a part in the formation of the banded structure by favourably influencing the

112

Hatfield Memorial Lectures Vol. II 950 qOO

J!.

o

w"

~ 850 !;;:

~

~ 800

~

750

Fig. 12

Influence of tin on the positions of the AC1 and AC3 points. QC REGION 20 100 200 300 400 500

bOO

700

~-----.-------~I~~~~-----'----~i Amplification

21SF

bOO

Fig. 13

700 )' REGION

800

QOOOC

Differential dilatometry of tin bearing steels. Reference specimen (I), O.035%Sn; comparison specimen (III), 0.261%Sn.

separation of the ferrite; this suggestion was later taken up by Thompson and Willows. 10 It might appear that oxygen should have a very considerable effect on the ability to form segregations, since Chipman's coefficient 1 - k equals 0.98; it has, however, no effect on the A3 transformation temperature and only a very limited solubility in iron. It is therefore quite probable that this element has only a very small effect, largely through the action of oxide inclusions. The alloying elements nickel, chromium and molybdenum, and also silicon, manganese and copper, suffer much less segregation than phosphorus, arsenic and tin. Nickel and manganese reduce the temperature of the A3 point and thus partly counteract the

The Mechanism of Formation of Banded Structures

113

effect of the impurities discussed before. This may provide an explanation for the fact that, in certain special steels (which are also generally much freer from impurities), the annealed structures sometimes differ from those usually found in carbon steel. Since metallic elements can readily be added in relatively large amounts, their segregation, although small in proportion, may nevertheless be appreciable in absolute terms. Various authors have attempted to explain the production of certain banded structures by manganese39,40 which lowers the temperature of the A3 point and thus delays the formation of proeutectoid ferrite during cooling; the ferrite forms preferably in bands having a low manganese content. A structure of this type has been reported by Schwartzbart+' in a steel containing 0.21 %C and 1.47%Mn; slow cooling gave rise. to the formation of pearlite bands in regions rich in manganese, and ferrite bands in areas with a low manganese content. Inclusions in steel which have been arranged into laminae through plastic deformation can, at least theoretically, play a part in the formation of the banded structure, either by favouring the physicochemical reactions at the interfaces, or by locally modifying the chemical composition of the steel and thus altering its transformation points if they are slightly soluble in iron at a high temperature. Any theories which ascribe to the inclusions an essential role in the formation of the banded structure+' are, however, certainly incorrect; for a given steel and for different cooling rates, inclusions of a particular type may be found sometimes in the ferrite and sometimes in the pearlite. Other reasons for rejecting such theories have already been suggested,10,13 among them the effect of a homogenisation treatment at a high temperature which partly destroys the banded structure without affecting the inclusions. We can therefore coincide that the inclusions have very little or no influence on the formation of banded structures.

CONCLUSIONS The considerations outlined above lead to the conclusion that the banded structure in steel is produced by chemical heterogeneity due to dendritic or small scale segregation during the cooling of the ingot. Each element which in solid solution produces a displacement of the temperature of the upper transformation point A3 of the steel, accompanied in some cases by a horizontal displacement of the transformation curve (Fig. 9), can give rise to a banded structure in accordance with the conditions and the mechanism which we have attempted to describe. Phosphorus plays an important part in the case of carbon and low alloy steels and with bands of proeutectoid ferrite, while other impurities, especially arsenic and tin, can augment this effect. In general, and this applies especially in the case of alloy steels, most of the clements which can produce dendritic segregation influence the mechanism, either by reinforcing or by counteracting the effect of the impurities phosphorus, arsenic and tin. The action of carbon is of a special type. It diffuses so rapidly that its content can be assumed to be practically uniform throughout the metal in the austenitic state. During

114

Hatfield Memorial Lectures Vol. II

cooling through the transformation points, the effect of the segregated elements is either to reconstitute the original segregation of the carbon at the instant of solidification, or, on the contrary, to reverse this segregation. During certain transformations the increase in the carbon content of austenite which has not yet been transformed may be so marked that the final heterogeneity shown by the micrograph can be much more pronounced than that caused by the segregation proper. A more detailed understanding of the mechanism of the formation of the banded structure can be very useful in increasing our knowledge of dendritic segregation and its results, in particular since some of the phenomena investigated are capable of very interesting generalisation. For example, the ghost lines of large forging ingots often exhibit a migration of the carbon under the influence of segregated elements, in particular phosphorus, on the scale of the ghosts, which is similar to that shown on the dendritic scale by the banded structure. Dr Hatfield always showed a particular interest in the heterogeneity of the transformations of steel, and he made an important contribution to our knowledge in this field. I have therefore been very happy to present this Lecture in honour of his memory.

REFERENCES 1. P. BASTIEN: 'La formation de la fibre dans les produits metallurgiques corroyes', Semaine de l'Etude de la Physique des Metaux, 1950. 2. Committee on the Heterogeneity of Steel Ingots: Iron and Steel Institute Special Report, 19261939,1-9. 3. ]. STEAD:]. Iron Steel Inst., 1915, (1), 140p. 4. H. HOWE: The Metallography of Steel and Cast Iron, New York, 1916. 5. A. SAUVEUR:Metallographie dufer et de l'acier; Gauthier Villars, Paris, 1937. 6. ]. WHITELEY:]. Iron Steel Inst., 1926, 213p. 7. P. CATTIER, C. DUBOIS,]. BLETON and P. BASTIEN:Rev. u«, 1950,47,619. 8. P. CATTIER, C. DUBOIS,]. BLETON and P. BASTIEN:Rev. u«, 1953,50,276. 9. C. BENEDICKSand H. LOFQUIST:Non-metallic Inclusions in Iron and Steel, Chapman and Hall Ltd, London, 1930. 9a.A. PORTEVIN: Rev. u«, 1913, 10,88-90. 10. F. C. THOMPSON and R. WILLOWS:]. Iron Steel Inst., 1931, (II), 151p. 11. Association Technique de la Siderurgie Francaise: Minutes of the Commission des Ingenieurs de traitements thermiques. Meetings of29.5.51, 21 and 22.11.51,20 and 21.5.52,20 and 21.1.53. 12. ]. LAVENDERand F.]ONES:]. Iron Steel Inst., 1949,163,14. 13. H. SCHWARTZBART:Trans. Am. Soc. Met., 1952,44,846. 14. G. REMY: Technical Information Circular of the Centre de Documentation de la Siderurgie Francaise, Vol. 9, 1952, 682. 15. C. ]ATCZAK, D. GIRARDIN and E. ROWLAND: Trans. Am. Soc. Met., 1956,48,279. 16. H. FINNISTON and T. FEARNEHOUGH:]. Iron Steel Inst., 1951, 169,6; discussion: 170,21. 17. J. CONVEY and]. OLDFIELD:]. Iron Steel Inst., 1946, (2), 473p.

The Mechanism of Pormation of Banded Structures 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41.

J. TRILLAT:

115

Rev. Sci., 1940, 4, 212; Metall. Rev., 1956, 1, 3.

u«,

G. HOMES and]. Gouzou: Rev. 1951,48,251. K. WOLFE and I. ROBINSON: Met. Treat. Drop Forging, 1949-50, 16,209. W. BETTERIDGE and R. SHARPE:]. A. KOHN: Rev. 1954,51,503.

u«,

Iron Steel Inst., 1948, 158, 185; discussion:

160,29.

R. CASTAING: Thesis, Bulletin Bimestriel de l'Office National d'Etudes et de Recherches Aeronauti-

ques, 1956, (55). j. PHILIBERT and C. CRUSSARD: Rev. u«, 1956,53,461. B. LARSEN: Trans. Am. Inst. Min. Met. Eng., 1945, 162,413. ]. Fuss: Bull. Cercle Etudes u«, 1943,3,341. A. KOHN and]. DOUMERC: Rev. Met., 1955,52,249. K. LEWIS: Iron Coal Trades Rev., 1955, 170, 125, 189,291,309,381,437; A. KOHN: Rev. u«, 1953,50, 139. D. KASANORSKY: Izvest. Akad. Nauk. SSSR., 1955,100, 1073. K. WOLFE and I. ROBINSON: Iron Coal Trades Rev., 1949, 158,337. P. BJORNSON and H. NATHoRsT:]ernkontorets Ann., 1955, 139,412. M. RE,y: Rev. u«; 1951,48, 167. M. GROSSMANN: Trans. Am. Inst. Min. Met. Eng., 1942,150,226. A. PORTEVIN and P. CHEVENARD: Compt. Rend., 1923,176,296. C. ROQUES, C. DUBOIS and P. BASTIEN: Autumn Meeting of the Societe Francaise de Metallurgic, 1956. D. WALKER: Australasian Eng., 1953, May, 63. ]. CHIPMAN: Chapter 16, Basic Open Hearth Steelmaking, AIME, 1951, New York. H. LE CHATELIER and B. BOGITCH: Compt. Rend., 1918, 167, 472; Rev. u«, 1919, 16, 129. B. LARSEN: Bureau of Mines, Technical Paper No. 466, 1930. A. PHOMOROFF: Metallurgia, 1936,13,179. N. ZIEGLER: Rev. u«, 1911,8,655.

TWELFTH

HATFIELD

MEMORIAL

LECTURE

Phenomena Occurring in the Quenching and Tempering of Steels G.

v.

Kurdjumov

At the time the lecture was given Academician Kurdjumov was the director of the Institute of Metallography and Metal Physics, Central Research Institute for Ferrous Metallurgy, Moscow. The lecture was presented at the Firth Hall, Sheffield University, on 25 November 1959.

I wish to expressmy thanks for the honour you have extended to me in inviting me to deliver the Twelfth Hatfield Memorial Lecture. I know that having chosen for this lecture such a wide subject as the quenching and tempering of steels I have taken upon myself a very difficult and complicated task. However, I am encouraged by the fact that, although many problems remain, both the general nature and many details of the quenching and tempering processes can now be described in terms of our general physical ideas of the phenomena occurring in solids.

Problems associated with quenching, such as the nature of martensite, the reasons for the high hardness of quenched steel, and the nature of the austenite-martensite transformation used to be considered the most complicated and uncertain in metallography. From Sauveur's report in 1926 on 'The current theories of the hardening of steel' ,1 thirty years later one can see how contradictory were the ideas of metallurgists about these phenomena. But in the three decades that have elapsed since that time the situation has been completely changed. The decisive factor in this change has been the application of X-ray structural analysis. This has enabled the crystal structure of quenched steel and its detailed changes at different stages of tempering to be determined, and has made possible the study of the atomic mechanism of the phase transformation on quenching and tempering. In choosing this subject for the lecture I had also in mind a further purpose. In the USSR problems of quenching and tempering of steels have been the subject of many investigations by physicists, and although some progress has been made in these investigations, many of them are not known in the UK and other countries. Therefore while giving a general account of the quenching and tempering of steels, at the same time I wish to outline some of the results obtained in the laboratories with which I have been connected. To my colleagues working in the same fields I apologise for not acknowledging here all those who have made valuable contributions to the solution of the problems under investigation.

117

118

Hatfield Memorial Lectures VoL II QUENCHED STEEL

The Nature of Martensite The special properties of steel obtained as a result of quenching are determined by the appearance in the cooling process of martensite, the main structural constituent. In each grain of initial austenite a large number of martensite crystals develop, usually having the form of plates. The thickness of these plates is less by at least an order of magnitude than their other dimensions. The size of the martensite crystals depends upon the state of the austenite grains, and the more uniform and perfect the crystal structure of the austenite, the larger are the martensite crystals. If the austenite grain structure is very non-uniform and there are imperfections in the crystal structure, the martensite crystals can be very small. The martensite has a body centred tetragonal lattice not very different from that of the cubic a-iron, the difference between them depending on the carbon content of the steel. 2-4 The important results on the effect of quenching conditions on the lattice parameters of martensite=f are as follows.

(1) If the quenching is carried out from the single phase austenite field and the cooling rate is sufficiently high to prevent the decomposition of austenite and the tempering of martensite during the quench, then the lattice constants a and c are independent of the quenching temperature and cooling rate, and depend only on the carbon content. The dependence is: c

= ao + = ao -

0.0118 P 0.015 P (c/ a) = 1 + 0.0467 P

a

and ao = 2.861 A, is the lattice constant of a-iron. 8 (2) The lattice constants do not depend on the carbon content if the quenching carried out from the same temperature in a heterogeneous field.

where p

= wt-%C,

is

From these experimental data it follows that (i) (ii)

martensite is a supersaturated solid solution of carbon in a-iron the martensite lattice constants are determined by the carbon content of the initial austenite (iii) martensite constains in solid solution, the same quantity of carbon as is dissolved in the initial austenite: transformation of austenite into martensite is a diffusionless process. It occurs without change of concentration of the solid solution, and only a change of lattice results.

In martensite as in austenite the carbon atoms are distributed in the interstices of the lattice. They are located at [OOY2] and [Y2 Y2zero] between iron atoms in the direction of the tetragonal axis:" carbon atoms are distributed in these interstices at random. The

Phenomena Occurring in the Quenching and Tempering of Steels martensite lattice has a definite orientation expressed by: (Oll)m [111]m

relationship

119

to the austenite," which can be

II (111)A II [101]A

Peculiarities of the Structure of Quenched Low Carbon Steel In carbon steels containing less than O.6%C, partial tempering of martensite occurs even with the most drastic quenching." Untempered martensite is, however, readily obtained in hypereutectoid steels. As the carbon content is lowered, the cooling rate must be steadily increased to obtain undecomposed martensite. This is because the martensite start point (Ms) increases with decreasing carbon content, while the rate of martensite decomposition increases quickly with temperature. The degree of decomposition for various martensite crystals is different because the transformation extends over a large temperature range. As a result the tetragonal structure becomes non-uniform.4,5,8,10 Undecomposed martensite with an axial ratio depending on the carbon content as in plain carbon steels (Fig. 1), can be obtained in steels having less than 0.6%C by adding alloying elements which lower Ms. For example, in a steel with 0.S3%C and 1.9%Mn undecomposed martensite can be obtained with the doublets arising from the tetragonality separated on an Xi-rayphotograph."! By introducing a suitable quantity of manganese and nickel, undecomposed martensite is obtained on quenching steels with O.S-0.2%C.12 Table 1 shows the results of the determination of c/ a. Undecomposed martensite was recently observed by the same method in steels containing 0.32 and 0.44%C and about 19%Ni.108

/

'·06

~

/~

1-04

~

u

0

t=

-c a:: 1·02

V

V

0·2

Fig. 1

Dependence

Q. 6

V

~

/'

0'4

/

/

•

Extracted martensite powder (ref~. 22,24)

.to Alloy steels with lowMs point (ref. 12)

0-8

C ,0/0

'·0

'-2

1-4

1·6

of the axial ratio c/ a of the tetragonal lattice of martensite upon carbon

content.

120

Hatfield Memorial Lectures Vol. II

Table 1 Axial ratio of the tetragonal lattice of martensite in quenched alloy steels with carbon contents less than 0.6%12 Content

0/0

C

Mn

Ni

Cu

Ms

cia

7.77 6.71 6.40 6.20 5.67

3.86 3.06 3.45 1.50 0.27

2.64 3.00 3.00 3.30 3.35

65 70 50 25 5

1.012 1.016 1.018 1.023 1.026

0.25 0.31 0.35 0.44 0.57

However, the question of the martensite crystal structure in steels containing 0.2-O.S%C cannot be considered as solved. By using single crystals9,13 we were able to obtain the reflection (002) in the absence of reflection (200), and vice versa. It was found that in steels containing 0.2-Q.S%C the reflection (002) was divided even when the possibility of decomposition was excluded owing to the low Ms.12 This can be explained by supposing that some martensite crystals are tetragonal and some cubic. From the ratio of the intensity of the spots it follows that the quantity of crystals with the tetragonal lattice is not large at 0.2%C, but it increases with the carbon content in the range 0.2o/0-0.S%C. The existence of two lattices of the martensitic phase is well known in the B' and 1phases of alloys of Cu-AI or Cu-Sn.35 The question of the simultaneous presence of cubic and tetragonal lattices in the supersaturated solid solution of carbon in a-iron in low carbon steels is still being investigated. The rate of martensite decomposition is determined not only by the mobility of the atoms, which quickly increases with temperature, but also by the driving force for decomposition, which becomes less when the supersaturation of the solid solution decreases. Therefore in plain carbon steels with about 0.1 %C and less, although M, is relatively high (about SOOOe), undecomposed martensite can be obtained provided the quench is sufficiently drastic. This is shown by the dependence on carbon content both of the breadth of X-ray diffraction lines, and of the hardness (Fig. 2) .101 Similar relationships have been observed in low carbon steels containing manganese or chromium (Fig. 3). Undecomposed martensite of about 0.1 %C can be obtained much more easily by adding elements raising the stability of the a solid solution. 14,15 For example, in steels containing titanium, vanadium or molybdenum the decomposition of martensite on tempering shifts to relatively high temperatures (4SQ-SOO°C), (Fig. 28). Obviously, to obtain undecomposed martensite on quenching such steels it is not necessary to have very drastic cooling in the temperature range below Ms' nor is it necessary to add elements that decrease ~. For these steels the difficulty is to create conditions for supercooling austenite to Ms' i.e. for preventing the precipitation of ferrite, or for preventing normal Y-7a transformation at high temperatures. These processes may be prevented either by a sufficiently high cooling rate in the temperature range A3 to Ms' or by introducing elements preventing Y-7CX transformation with normal kinetics.lv-P' The rate of normal Y-7a transformation can be also decreased to a great extent by removing (or making ineffective) the imperfections of the crystal structure of austenite induced by the transformation on heating, and which promote the formation of nuclei of the a phase on subsequent cooling. These imperfections may be

Phenomena

Occurring in the Quenching

and Tempering of Steels

__--~--~----~--_r--~._--~--~50U

121

V') V')

w

40~

z

« ::c

4·2 J8 0

E E

~

~3·4

:r:

o«

w tX:

co

w

Z

:::; u

12 I" w

~ o SlJ... W

>

~

4W

8

OL----L---O~OL4--~---O-LO-8---L---O~.-12--~O C,O/o

Fig. 2

Hardness (Re) , line breadth (~211) and coercive force (H) in low carbon steels quenched from 100Qoe into iced water, as a function of carbon content.lv! MANGANESE STEELS,2·0-2·Zo/oMn 50 •••••c)

5'0

Rc

...... ~t'

_/ 17

..J7

,0

/a

B

Vc

~

V

40~ VI

30~

Z -e :r:

20

~

24~ ",It

0

~o

Me

2·6

,,'

.~~-t ,./

.,,'

~ 0-02

-

"",.

0·04 CARBON

Fig.3

,P 10- •••••

oY

4·6 4·2

/

0-06

0·08

CONTENT,

0-10

O'IZ

8 0-14

%

Hardness (R), line breadth (P211) and coercive force (H) of alloy steels containing

2%Mn as a function of carbon content (101). B

= ~, Me = R;

122

Hatfield Memorial Lectures Vol. II

removed (or made ineffective) by heating austenite to a sufficiently high temperature.I? Thus by quenching from 1150°C and higher, we were able ·to prevent the formation of 'normal' a-iron nuclei and make the y~a transformation proceed martensitically in iron containing less than O.Ol%C.2o The iron quenched in this way showed the same hardness and line breadth as after heavy cold work (Fig. 4).

(\

Q

V - -- -~

1\

~V \

r-,

,

6000 Hv-SO

j

IObO°

Hv·120

Fig. 4 Hardness (HV) and intensity distribution in Ka1-Ka2 lines of (220) reflections (Fe radiation) of a-iron containing about 0.01%C, after different heat treatments. Iron drastically quenched after heating above 1150°C has a martensitic microstructure; the hardness and line breadth are almost the same as for cold worked iron.

The lattice of pure iron martensite is obviously the same as that of annealed iron, and crystals of the martensite differ from the annealed material only because of the rnicro- and subrnicrostructure (distortions of crystal structure). The difference in structure is caused by the difference in the mechanism of formation. The martensitic mechanism of transformation leads to a state similar to that obtained as a result of cold work.2o,26 In the presence of carbon the martensite crystals differ from annealed a-iron not only in physical state, but also because of the presence of carbon in solid solution. Up to a certain content of dissolved carbon the martensite lattice remains cubic. In this, carbon atoms are randomly distributed in the lattice between iron atoms along the three cube axes. The martensite lattice probably remains cubic even with 0.1%C, although this could not be determined by direct measurements of lattice constants. The lattice is distinctly tetragonal with a content of O.2%C, 12 so that it seems that regularity in the distribution of carbon atoms and so the tetragonal synunetry of the martensite lattice, appears between 0.1 % and O.2%C. Imperfections in the Structure of Martensite Crystals The diffuseness of lines and their low intensity in X-ray patterns of quenched steels shows that martensite crystals have a very imperfect structure. In investigating these peculiarities

Phenomena Occurring in the Quenching and Tempering of Steels

123

of the X-ray diffraction picture the use of austenite single crystals'> and of electrolytically extracted martensite powder22-24 have played an important role. With the former, a regular orientation relation between the martensite and austenite lattices makes it possible to avoid difficulties connected with overlapping of the doublet components due to tetragonality and to select for investigation more favourable rcficctions.F! The study of extracted martensite powder showed that the isolated martensite crystals have a tetragonal lattice with the same dependence of lattice constants on carbon concentration as in quenched specimens, but that the diffraction lines of the martensite powder are much sharper (Fig. 5). The first fact shows that the lattice constants of martensite are determined only by the presence of dissolved carbon, and there is no connexion whatsoever with the stresses. The increased resolution of the X-ray pattern enabled us, first, to understand the nature of the microstresses causing considerable line broadening, second, to obtain more reliable data on the dimensions of coherent regions in martensite crystals and third. to make precise measurements of the intensity of lines. 23,25

Low 6

Fig.5

X-ray photograph of martensite powder electrolytically extracted from a quenched steel containing about 1% carbon.F

Regions if coherent X-ray scattering and the elastic diformationsofmartensite (m icrostresses)

crystals

The angular dependence of line broadening in martensite showed that the reasons for diffuseness are (a) the presence of microstresses (non-uniform elastic deformations of microregions) and (b) the small size of the regions of coherent X-ray scattering.F' The latter, within the precision of measurement, does not depend on the carbon content and is about equal to 200-300 A both for carbon free martensite-v-e? and high carbon martensite.21,23 The dimensions of these regions can be estimated more reliably by measuring the line breadth of extracted martensite powder. In this case the broadening is determined almost entirely by the smaIl size of these regions and it is proportional to the secant of the angle of reflection. 23 When the martensite crystals are separated, that part of the line broadening which is proportional to the tangent of the angle of reflection disappears. For steels with high carbon content this part may be quite large. For example, the breadth of the (220) line of a bulk specimen of quenched steel is greater than 100 mrad.s! but after separation of the martensite crystals this breadth falls to 20-30 mrad (Fe radiationj.P From these data it was concluded that the line broadening which is proportional to the tangent of the angle of reflection is determined by the non-uniform elastic deformation of

the martensite crystals (e.g. the bending of martensite plates) caused by the mutual

124

Hatfield Memorial Lectures VoL II

constraints exerted by the crystals on one another. On releasing martensite crystals from their surroundings the forces causing their elastic deformation disappear.i" The magnitude of the elastic deformation /1a/ a of a martensite crystal in quenched steel increases greatly with increasing carbon content. In a steel containing 0.1%C it is between about 2.5 X 10-3 and 3 X 10-3, i.e. it is considerably higher than the elastic deformation of microregions in cold worked iron (-1 X 10-3). With high carbon content the elastic deformation may increase to about 10-2. Thus, the small size of the coherent regions is one of the internal characteristics of martensite crystals; it is the same whether a martensite crystal is in a bulk specimen of quenched steel or separated from it. But the non-uniform elastic deformations are not themselves inherent characteristics of the internal structure. They arise because of mutual constraints caused by the shape changes of microregions during the transformation, and they disappear when the martensite crystals are extracted. However, the maximum value of the elastic deformation of martensite crystals of a given steel is a measure of the mechanical properties of these crystals, first of all of their elastic limit. The analysis of line broadening thus shows that the elastic limit of martensite crystals increases rapidly with increasing carbon content. Since the thickness of martensite plates is of order of 10-4 cm, there are about 100 coherent regions in the plate thickness, so that these regions are deformed practically uniformly. 25

Static displacements and thermal vibrations The presence of carbon atoms in the interstices of the iron lattice leads to a displacement of the iron atoms from their ideal positions. Obviously this displacement must be largest for the atoms which are nearest neighbours of carbon atoms. Also the displacements in the direction of the tetragonal axis must be considerably larger than those in perpendicular directions.f because of the difference in average distance between iron atoms along the c and a axes revealed by the dependence of the lattice constants on the carbon content. This behaviour is in accordance with the conj ectured co-ordinates of carbon atoms in the martensite lattice; the distance between carbon and iron atoms along the c axis is considerably less than that in the perpendicular plane. Determination of the mean square displacement of the iron atoms from their ideal positions by measuring intensities confirms this supposition.P'v'? According to data obtained from separated martensite powder (Fig. 6) the mean square displacement along the c axis appears to be twice as great as that along the a axis.23 Figure 7 shows the variation with carbon content of the mean square displacement of iron atoms (not taking in account anisotropy), obtained by measuring intensity using molybdenum radiation.I? The presence of carbon atoms dissolved in the a-iron leads not only to a static displacement of atoms in the martensite lattice but also changes the interatomic forces. Thus the mean square thermal displacement of atoms increases with the carbon content at a given temperature (Table 2).30 Figure 8 shows the results of determinations of the Debye characteristic temperature from data on the intensity of martensite lines at room temperature and at the temperature

Phenomena Occurring in the Quenching and Tempering of Steels

125

Fig. 6 Relative intensity (II P) (11101 PllO) of reflections of a-iron and extracted rnartensitic powder (a'). P is the multiplicity factor. Note the weakening of the martensite reflections, especially of those with a higher third index.F' 12 10

./

.. V

~ o

?

/1

0

4

V

1

0'2

0-4

0-6 C,%

o-e

,·0

"Z

Fig. 7

Mean square static displacements of iron atoms in the martensite function of carbon content (using Mo radiationj.P?

Table 2

Mean square dynamic and static.atomic displacements in martensite lattices with different carbon content

lattice as a

Carbon content ~, 0/0

8, oK

-fL:i2din' A

..JU2st' A

a-iron 0.08 0.10 0.35 0.84 1.0

430 435 435 390 365 360

0.115 0.114 0.114 0.126 0.133 0.136

0.045 0.057 0.059 0.085 0.100

of liquid nitrogen. In accordance with this, Young's modulus for quenched creases with increasing carbon content (Fig. 8, upper curve).31

steel de-

NATURE OF THE HIGH HARDNESS OF QUENCHED STEEL There

are two kinds of structural

the martensitic

mechanism

features associated with quenched

of growth

results in a particular

steel. 25 First,

fine micro-

and

126

Hatfield Memorial Lectures Vol. II fl

E

T

0/0

~fr-+---~--+---~--4---~O

3S0'"--"'----'--_...a..-_~_.......&__~

o

0·2

0·4

0·6

CA RBON CONTENT,

O-S

0/0

J-O

1-2

Fig. 8 (a) Debye characteristic temperature (8) of quenched steels calculated from X-ray photographs obtained at +20°C and -194°C with Mo radiation.J? (b) relative change in Young's modulus (~E/E) of quenched and tempered steel.U submicrostructures; a large number of martensite plates lie within each austenite grain (in low carbon steel austenite transforms almost completely), and there are coherent regions of submicroscopic size inside the plates. There is also a non-uniform elastic deformation of the plates and a regular orientation of martensite to austenite. Second, owing to the diffusionless character of the martensite transformation the martensite crystals consist of a supersaturated solid solution of carbon in rz-iron existing only in a metastable state. Structural changes of the first kind, which are not dependent on the presence of carbon, have the same effect on mechanical properties as cold plastic deformation. Thus, the increase in strength of quenched pure iron or ferrous alloys (if the y~a transformation occurs as a purely martensitic one) is nearly the same as that caused in these materials after heavy cold work. Quenching carbon free iron, like cold working, leads to a different micro- and submicrostructure compared with annealed a-iron, but not to any difference in chemical composition. The mechanical properties of the lattice itself are therefore those of the carbon free annealed a-phase. The situation is quite different when steels are quenched, for this not only leads to a fine micro- and submicrostructure but also alters the properties of crystals themselves in small regions, because of the formation of a supersaturated solid solution of carbon in the a-phase. Owing to the carbon in solution the limit of elastic deformation rises rapidly with increasing carbon content, and this high elastic limit of the martensite crystals is a fundamental cause of the high hardness of quenched steel. Thus the quenched steel possesses considerably higher hardness than cold worked iron, or iron quenched to obtain martensite. The influence of the properties of martensite crystals on the strength can be clearly seen by considering the properties of quenched low carbon steels. Figure 9 shows the

Phenomena Occurring in the Quenching and Tempering of Steels

127

relation between the hardness of quenched low carbon (0.03-0.12%C) steels and the value of the elastic deformation in microregions (microstrains) obtained from the measurements of line broadening.101,33 This proportionality between hardness and microstrains might be treated as an influence of microstrains themselves on hardness. But many observations suggest that such properties as yield point and hardness remain the same while microstrains vary within wide limits, if the micro- and submicrostructure do not change. A special investigation confirmed this point (Fig. 10).32 Values of the microstrain da / a determined from measurements of line breadth, can be considered as characteristic of the elastic limit of crystals in microregions. The above mentioned increase of the hardness with J1a / a (Fig. 9) is therefore due to the increase in the strength of the crystals in microregions, but not due to the presence of microstrains J1a / a themselves.25,26,33 These ideas on the importance of the effect of the elastic limit of crystals in microregions on the strength of metals (after martensitic transformation or cold work) were confirmed by an investigation of iron-silicon alloys.Y'Figure 11 shows that the hardness and microstrains Lla / a of cold worked alloysincrease with silicon content, in parallelwith a rise in the hardness of annealed alloys. Thus the increase in the strength of steel as a result of quenching is determined first by the formation of a fine micro- and submicrostructure, and second, by the high elastic limit of the martensite crystals themselves, associated with the presence of dissolved carbon. The first factor makes a contribution which depends slightly upon carbon content. The second factor makes a contribution to the strength of quenched steel which increases with the carbon content, and is the main cause of its high hardness.F>

3'3

~Id

~ 2·, ~--......---t-I----iI--------iI------f----1 Z :< c: t;;

~ 1'7~--~---"""---~~------i1-------f------1 u

>:

Fig. 9

Relationship between the hardness of quenched low carbon steel and the magnitude of the microstrains L1a / a.101,103

128

Hatfield Memorial Lectures VoL II 85E E

751"

280

on

5 E

..au

30

~

2'8 2·4

oj

<)0

1·6 1'2

0·8 0'4 0

Fig. 10 martensite

Changes in characteristics of fine structure and in the strength ofFe-Ni (25%Ni) after heating at different temperatures.V O's is yield stress, D is size of regions of coherent X-ray scattering .

......---~--_r------r----r-----,6-0 B ..£)

----~----+---

Fig. 11

Q >C

4'0 0

Changes in the hardness and microstrains of cold worked Fe-Si alloys, and of the hardness of annealed alloys, as a function ofSi content, D is size of regions of coherent X-ray scattering.

Phenomena Occurring in the Quenching and Tempering of Steels NATURE OF THE AUSTENITE-MARTENSITE

129

TRANSFORMATION

If steel is rapidly cooled from a high temperature the decomposition of austenite by diffusion process is prevented, However, austenite in plain carbon steel cannot be completely supercooled to room temperature because the martensite transformation intervenes. This transformation of austenite to martensite (A---7M) has a number of peculiarities sharply distinguishing it from other types of transformation in the solid state.34-37 The A---7M transformation was the first one in which these peculiarities were found, but we now know that transformations having particular characteristics of the A---7M transformation take place in many metals and alloys, and that this type oftransformation is predominant in temperature ranges where diffusion is slow. All transformations of this type are now called martensitic. The main characteristic of martensitic transformations is a special mode of formation of the crystals of the new phase, consisting of a co-operative regular rearrangement of atoms in which the relative displacements of neighbouring atoms do not exceed the interatomic distance; and as a result of the rearrangement of the atoms a macroscopic shear is produced.35,38 This shear is revealed by the relief produced on a polished surface by the transformation. The co-operative and regular character of the atomic rearrangement makes it possible for the transformation to proceed in a temperature range where diffusion is very slow. The shearing character of the lattice rearrangement causes large elastic deformations during the growth of martensite crystals and it suggests, among other peculiarities, that stresses will be important in the kinetics of martensitic transformations. Metallurgists have for a long time been very interested in the nature of the A---7M transformation, not only because this transformation is connected with the important problem of the hardening of steel by quenching, but also because of its unusual character. To understand the transformation it is necessary to make clear the nature of the driving forces and the mechanism of formation of martensite crystals. We have also to find out the causes of the following main features of the kinetics: (a) the high rate of formation of martensite crystals and the absence of their further growth, and (b) the rapid damping of the process of formation of new martensite crystals when cooling is stopped, and the fact that the transformation extends over a temperature range. We have also to explain such other phenomena as the stabilisation of austenite, the anomalous influence of the cooling rate on transformation, and the formation of martensite by plastic deformation.

Factors Causing the Transformation The martensitic transformation of austenite occurs, as do all other phase transformations in the solid state, because below a definite temperature there exists a structural state of the alloy with lower free energy. Since martensitic transformation proceeds without change of concentration, it can be treated as a transformation in a one component system.V

130

Hatfield Memorial Lectures VoL II

Austenite and martensite are crystal modifications of the solid solution similar to polymorphic modifications of pure metals or chemical compounds.f The free energy of each of these modifications (FA and FM) is temperature dependent, and the intersection point of the free energy curves gives the temperature To of metastable equilibrium between austenite and martensite. This point is determined by the chemical composition of the solid solution and can be calculated if the thermal properties of both modifications of the solid solution are known.39,4o If the solid solution does not decompose by diffusion a driving force arises below To which makes the transformation possible, and the greater the free energy difference between austenite and martensite (FA-FM) the higher the driving force. However, transformation does not start immediately below To' but only at M, which is considerably lower (about 200°C in steels), (Fig. 12). When martensite is heated above To the reverse transformation must be expected.

CONCENTRATION

Fig. 12 Schematic diagram of 'reversible' martensite transformation. To is temperature of metastable equilibrium between the two crystal modifications of the solid solution; 1\11s is start of martensitic transformation on cooling; As is start of reverse martensitic transformation on heating. The reverse transformation of the martensite phase directly to that which is stable at high temperature, was first observed on heating the ~'-phase in Cu-AI alloYS.41,42 Favourable conditions here are a small rate of decomposition of the solid solution and a comparatively small hysteresis of transformation. The martensitic character of the reverse transformation was proved later.43,44 Owing to the high rate of decomposition of the supersaturated solid solution of carbon in a-iron, it is very difficult to observe the reverse transformation in steel

Phenomena Occurring in the Quenching and Tempering of Steels

131

even at extremely high rates of heating. However, the martensitic a~"{transformation on heating can be easily obtained in ferrous alloys. Both on heating and cooling, martensitic transformation leads to relief on a polished surface (Figs. 13 and 14).

Fig. 13 'Reversible' martensitic transformation in an Fe-30%Ni alloy.46 (a) relief arising from the martensitic Y-7(J.. transformation when a polished specimen was cooled in liquid nitrogen; (b) the same specimen repolished, heated to 600°C and quenched in water; the figure shows the relief caused by the reverse (J..-7Y transformation. The markings in the two photographs are reversed, hills becoming valleys and vice versa. Thus similar shears are involved in the two transformations but they are in the opposite sense.

Fig. 14

'Reversible' martensinc transformation in a Cu-24.8%Sn alloy."? (a) after quenching in water from 680°C, polishing and cooling in liquid nitrogen; relief cased by f3-7f3" transformation is shown; (b) same specimen after repolishing, heating for 5 s at 200°C

and quenching in water; relief is caused by the reverse

~"-7~

transformation.

132

Hatfield Memorial Lectures VoL II

The large hysteresis in the martensitic transformation makes experimental determination of To difficult. The main reason for this large hysteresis in a 'reversible' martensitic transformation is the energy of elastic deformation arising when the crystals of martensite form: extensive supercooling (or superheating) is necessary to make the free energy difference between both crystal modifications large enough to compensate this elastic energy. Martensite can be produced from austenite at temperatures above M, by plastic deformation, but the farther above M, the temperature of deformation, the greater is the amount of plastic deformation needed. At temperatures above Md, characteristic for a given alloy, it is impossible to produce martensite by means of plastic deformation.v' Our conception of the driving force of transformation implies that plastic deformation can stimulate the A---7M transformation only below To, i.e. in the temperature range in which martensite is thermodynamically more stable. At temperatures above To plastic deformation must cause the reverse M---7A transformation. Observations of A~M transformations caused by plastic deformation have been used for the experimental determination of To as a function of composition in Co-Ni and Pe--Ni alloys.40,49 Thus stresses playa big role in the A---7M transformation, but cannot by themselves be considered as the reason for it. They do not govern transformation but are an important factor influencing its kinetics.

Mechanism

of the A---7M Transformation

Formation of martensite crystals regarded as a rearrangement of atomic packing The martensitic transformation, involving the formation of crystals of a new phase inside those of the initial one, must proceed like all other processes of a similar kind, by means of the formation of nuclei and their subsequent growth. The experimental data about the A---7M transformation, particularly the analogy between martensite transformation and mechanical twinning.ff and the strictly regular orientation relationship," led to the above mentioned conclusion that the growth of martensite crystals consists of a regular diffusionless rearrangement of the lattice in which the relative displacements of neighbouring atoms do not exceed the interatomic distance.38,so,3s This conclusion was confirmed by the observation that ordered solid solutions remain ordered after martensitic transformations. From the mechanism of transformation the coordinates of atoms of both elements could be calculated and then verified experimentally (y-phase in Cu-AI and a'-phase in Cu-Zn).Sl-S3 The idea of cooperative regular rearrangement of atomic packing explains the high rate of formation of martensite crystals at low temperatures. This lattice rearrangement must take place in such a way that shears occur; as already mentioned, these are studied by investigating the microrelief caused on polished surfaces by transformation.46,47,35 The details of the atomic mechanism of formation of crystals of martensitic phases can be derived from the experimental observations of lattice relationship, habit planes and directions and amount of shear. 37

Phenomena Occurring in the Quenching and Tempering of Steels

133

Coherency between martensite and austenite lattices during growth One of the important assumptions about the character of the rearrangement of the atoms is the existence of coherency between the lattices of the growing martensite crystal and the austenite matrix.v> There is an order in the atomic arrangement on the boundary; atoms which were neighbours in the austenite lattice remain neighbours also on the boundary of the growing martensite crystals. Under such conditions, and with the shearing character of lattice rearrangement, high shear stresses must arise. These stresses increase with the growth of the martensite crystal. When the stresses reach some definite value, coherency is destroyed and the order of atomic arrangement on the austenitemartensite boundary disappears. The high rate of growth takes place only while the coherence is maintained, and is a consequence of the cooperative character of the atomic movements and the small size of the relative atomic displacements during the lattice rearrangement. The limited growth of martensite crystals can be explained by the loss of coherency-> since growth by means of non-ordered individual atom movements is not observable at low temperatures. The concept of coherent growth appears also to be useful in understanding the peculiar form of martensite crystals, for this must satisfy the condition that the elastic energy be a minimum for a given volume of martensite.

Thermoelastic equilibrium and (elastic) martensite crystals The idea of coherent growth led in the following way to the prediction and subsequent observation of the phenomena of thermoelastic equilibrium and 'elastic' martensite crystals. During the coherent growth of the martensite crystal a large elastic energy arises. Under certain conditions this positive part of the free energy change may increase more rapidly than the negative part, the free energy difference between the new and original crystal modifications. Thus the total free energy may pass through a minimum as the dimensions of the martensite crystal increase. If this happens before the loss of coherency, the growth will stop and the martensite crystal will be in thermoelastic equilibrium with the parent phase. Raising the temperature will then cause the crystal to shrink and lowering the temperature will cause it to grow.i" Such a phenomenon was observed in the martensitic transformation ~1 ~1 in alloys of Cu-Al with addition of nickel and manganese (Fig. 15), and subsequently in some other alloys.

Conditions for the Occurrence of 'Normal' and Martensitic Mechanisms of Transformation In steels and other eutectoid alloys the phases arising from martensitic transformation are metastable: for example, martensite in steel, the P' and 1-phases in Cu-Al alloys; the ~' and p"-phases in Cu-Sn alloys; and the a' and p'-phases in Cu-Zn alloys.P> In carbon free iron or ferrous alloys and in the metals cobalt, zirconium and titanium and their alloys, phases stable at low temperature are formed by martensitic transformation. Transformation of the high temperature modification to the low temperature one may occur in these systems either by 'normal' kinetics or by martensitic means: with the

134

Hatfield Memorial Lectures Vol. II

Fig. 15a Elastic martensite crystals in a Cu-Al-Ni alloy,102 (a) and (b) represent two places on the same specimen. Parts 1-3, cooling is from left to right; parts 4-6, heating is from left to right.

former the formation of the crystals of the new phase occurs by disordered atomic displacements as in recrystallisation; with the latter it occurs by ordered cooperative movements of atoms. The first is possible when the equilibrium temperature lies in the temperature range where diffusion is sufficiently rapid: for example, in pure iron or some of its alloys. If the equilibrium point lies in a temperature range where diffusion is slow, the transformation proceeds martensitically. As diffusion can be prevented by rapid cooling, the transformation y~a in pure or alloyed iron may be made to proceed either by the first mechanism or by the second, by varying the cooling rate.16-18 On rapid cooling normal transformation is prevented and the transformation proceeds below M,

Phenomena Occurring in the Quenching and Tempering of Steels 135

Fig. 1Sh Elastic martensite crystals in a Cu-Al-Ni alloy.J02 (a) and (b) represent two places on the same specimen. Parts 1-4, cooling is from left to right; parts 5-8, heating is from left to right. (which lies well below To owing to the considerable hysteresis caused by the large elastic energy involved). In some ferrous alloys normal "{-7U transformation proceeds comparatively slowly. The temperature dependence of the rate of the normal "{-7U transformation could therefore be measured for a number of alloys of iron with chromium, nickel and other elernents.tv-"? For iron alloys with A3 about SOO°C the transformation rate at first increases with decreasing temperature, reaches a maximum at about 700°C, and then decreases and becomes extremely slow below 550°C (Fig. 16). In pure iron and iron alloyed with tungsten, molybdenum and cobalt, the rate of the normal Y-7U transformation is extremely high and the temperature dependence could not be measured. In these alloys it is difficult to prevent normal transformation and to obtain a martensitic structure on quenching. However, as already mentioned, in pure iron we were able to decrease considerably the rate of normal transformation by first heating the "{-iron to a high temperature. On subsequent drastic cooling the "{-7U transformation could be made

136

a:

Hatfield Memorial Lectures Vol. II

60~--~---4--+---4-

lJ.J V)

<{

I

0-

~ 40

!-+---I-+------,{---!--.."c.---t--"..o'

20

Fig. 16

25 TIME, min

30

45

Kinetic of'Y~a transformation in an iron alloy (7%Cr and 2%Ni); AC3 M,

=

=

50

830°C,

44SoC.17

to proceed martensitically.F" In titanium and zirconium either mechanism can take place depending on the cooling conditions, but as the equilibrium point in cobalt is low, only a martensitic mechanism occurs here. However, both mechanisms are possible in Co alloyed with elements which raise the temperature of equilibrium.

Temperature Dependence of the Transformation Rate Investigations of martensitic transformations in copper alloys led us to the conclusion that they can be considered as a process of nucleation and growth. It was further supposed that this process was not athermal and that at sufficiently low temperatures the transformation would proceed slowly.35 Studies of the kinetics of austenite-martensite transformation below room temperature showed that the rate of transformation did indeed depend on the temperature.v' In this range of temperature some phenomena which were considered as being characteristic of the martensite transformation are no longer present. It was established that: (i)

under these conditions transformation

can take place isothermally and continue for a

long time (ii) the transformation rate depends upon temperature and decreases with decreasing temperature (Fig. 17a). (iii) transformation can be completely suppressed by rapid cooling to liquid nitrogen temperature (iv) austenite supercooled to the temperature of liquid nitrogen transforms to martensite on heating up to room temperature, with a rate depending on temperature (Fig. 17b, c).

Phenomena Occurring in the Quenching and Tempering of Steels .fn

(IlV) ~n-

137

t",O (b)

o

0-

o

2

6

4

8

10

12

I

+ 100 !() -40-80

-fZO-f4O -160 -180 -190

TEMPERATURE,O(

14~6

-aoo

Fig. 17 Isothermal austenite-martensite transformation in a steel with O.7%C, 6.5%Mn, 2%Cu.54 (a) transformation at different temperatures; (b) logarithm of the initial rate of A----?M transformation versus the reciprocal of the absolute temperature; (c) transformation at -159°C after supercooling of austenite in liquid nitrogen. It was also found that even in this temperature range the following specific characteristics of martensite transformation are retained: (i) martensite crystals in steel grow with an extremely high rate even at low temperature (ii) the transformation extends over a temperature range: when cooling stops only a certain part of the austenite transforms, and isothermal transformation then proceeds only at a lower temperature. The temperature dependence of the rate of transformation is determined by the temperature dependence of the rate of nucleation, since the growth of each martensite crystal is limited, and proceeds rapidly within the entire temperature range; and the decrease of this transformation rate with decreasing temperature shows that thermal vibrations play an important role in forming martensite nuclei. The possibility of supercooling austenite to the temperature of liquid nitrogen shows that when the thermal energy is small, the formation of nuclei in a given specimen does not take place. However, transformation starts on heating as soon as the energy of thermal vibrations is sufficiently large: and the higher the temperature, the faster it proceeds.

138

Hatfield Memorial Lectures Vol. II

The temperature dependence of the nucleation rate is thus the same in the martensite as in all other types of phase transformation: the rate of the isothermal transformation below the martensite point at first rises with decreasing temperature, and then passes through a maximum (Figs. 18 and 19). However, if the martensite point is not sufficiently low there will be a temperature range in which the rate of isothermal transformation cannot be measured. The transformation rate is easily measured at temperatures below -SO°C. At higher temperatures (near room temperature) the energy of thermal vibrations becomes so large, that, when cooling is stopped, all nuclei of critical size which can be formed at a given temperature form in a short period of time. The formation of martensite nuclei is also slow at temperatures just below Ms' This is because the critical size of nuclei, and hence the free energy of nucleation, are large, so that large thermal fluctuations are necessary. When the temperature decreases (i.e. the supercooling increases) the critical size of nuclei decreases and the rate of nucleation rapidly increases. Therefore for normal steels the isothermal A~M transformation can be observed only either in the temperature range somewhat below M, or in the range below -SO°C. In a temperature range near room temperature isothermal transformation proceeds at a high rate (Fig. 19, showing a steel with M, = +lS5°C). To summarise then, at temperatures lower than -SO°C the small thermal energy causes a low rate of nucleation and permits isothermal transformation to be observed; this factor becomes more pronounced as the temperature decreases. In the temperature range near M, the large energy needed for nucleation makes it possible to observe the isothermal transformation; the effect of this factor decreases with increased supercooling. On lowering M, (by changing the steel composition) it is possible to displace the temperature range, where the energy of nucleation is large, to lower temperatures, so that it becomes possible to measure the rate of isothermal transformation also in the region near room temperature (Fig. 19, steel with M, = 8S °C). Thus in the kinetics of nucleation, thermal vibrations play an important role in the martensitic, as in other phase transformations. The high rate of nucleation at a temperature as low as room temperature, where diffusion processes are frozen, arises because the necessary atomic displacements for the formation of martensite nuclei formation require little activation energy. This is probably connected with the small relative atomic displacements required, and with their regular cooperative character. The low rate of nucleation and the high rate of growth of martensite crystals at temperatures considerably below room temperature, indicates that this growth requires an extremely small activation energy. Apparently the energy barriers become smaller during growth due to the increase in the stresses at the boundary of the growing crystal. 59 Therefore in the initial stages of growth of the martensite crystal the rate of growth is probably not so high as it becomes when the crystal is visible under the microscope.

Reasons Why Transformation Extends Over a Wide Temperature Range One of the main peculiarities of martensitic transformation is its extension over a wide temperature range; at any temperature in this range only a part of the original phase transforms into the new one. As the martensite crystals do not grow after reaching a

Phenomena Occurring in the Quenching and Tempering of Steels 139

200

160~---+-

20 UI °t'J1

u'

0

°01 1"")1 ,'••.. ,I 1 2:,

+1 I •..• 1 2:1 Q

-50

-100

aU:

TEMPERA1URE,oC

-150

~I

:,1

-£.

Fig. 18 Temperature dependence of the initial rate of isothermal A-7M transformation in Pe--Ni-Mn alloys with different M, points. (1) 22.7%Ni, 3.1%Mn, M, == +12°C; (2) 22.S%Ni, 3.4%Mn, M, == -30°C; (3) 23.8%Ni, 3.2%Mn, M, == -55°C . •. 20

I \

~

i'o

~z;

16

mV

0::

2 Z ~....

o

12

/ J'

V

~ w

>

~

-l

~

0

\ \

\, ~

}o -

+200

10

~, ..\~

\ \\

/

1 .•.'00

\ \ (III)

r.~..' \ (II)

01~.,f

8

i= -c

/

I

~~~

a

TEMPERATURE,OC

-lao

~ ~

-200

Fig. 19 Temperature dependence of the initial relative rate of the A---7Mtransformation in steels with different Ms points.f? (I) O.85%C, 2.2%Mn, Ms == +155°C; (II) O.95%C, 3.S%Mn, Ms == +85°C; (III) O.70%C, 6.S%Mn, 2%Cu, Ms == -SO°C.

140

Hatfield Memorial Lectures VoL II

certain size, the termination of transformation implies that the nucleation of the new phase stops while the original one is still present. The explanation of this phenomenon is one of the main problems of the theory of martensitic transformations. Two factors may cause the isothermal transformation to come to a halt: (i) the state of the retained austenite may change in such a way that the formation of new nuclei or their growth becomes more difficult. (ii) nuclei may not form homogeneously in the initial phase, but only in some sites where nucleation is easy. Factors of the first kind may be, for example, a rise of pressure causing a decrease of Ms; or an increase in structural imperfections hindering nucleus formation or growth. Factors of the second kind may be inhomogeneities in the austenite, such as frozen-in fluctuations of concentration of a dissolved element, or different kinds of imperfections of crystal structure causing local stresses which decrease the energy of nucleation; shear-stresses which are uniform in regions, whose size is comparable with that of the martensite crystals, etc. Factors of the first kind can hardly be important initially, because a small quantity of martensite is unlikely to change the state of the retained austenite significantly, but they can be important during the second half of the martensite transformation curve. Therefore, in the first half of this curve, factors of the second kind must play the main role. At a temperature just below Ms' those sites are used where, due to favourable variations of structure or concentration, the energy of nucleation is smallest. After these sites are used up, the process of nucleation will terminate because in all other places the energy of nucleation is too large to be provided by thermal fluctuations. At a slightly lower temperature less favourable places can be used for nucleation, and so on. For studying the influence of imperfections of crystal structure on the nucleation of martensite, investigations were carried out in which specimens were subjected to plastic deformation or neutron irradiation. Steels and ferrous alloys with an M, point lower than room temperature were used. Different amounts of plastic deformation were applied in the temperature range 20-200°C, and the effect of both kinds of preliminary working (plastic deformation and neutron irradiation) on the kinetics of transformation below room temperature were studied. The results of these investigations confirm the supposition that the martensite nuclei are formed not uniformly throughout the volume, but in sites where the structure is distorted. Both cold work and neutron irradiation can cause structural distortions in austenite which accelerate the formation of the martensite nuclei (Figs. 20-22). However, these distortions are very unstable and disappear slowly even at room temperature. It follows that the sites of easy nucleation are characterised by high local stresses in small volumes, for such stresses can partially relax at room temperature, and even below. However, after more extensive plastic deformation or neutron irradiation, and in the process of transformation itself, structural distortions of another kind arise in the austenite grains causing transformation to proceed more slowly, or even preventing it. These

Phenomena Occurring in the Quenching and Tempering of Steels 141

...>< i

r;::;---o 2

20~~---+----~--~--~

~216~4----+~~~--~--~

>1'"

<]<]

~

z

12t---+-----f-----+---t---t------1

o ~ Z

8t---+--+-+--~~;+-t---~

-c

e:

o

4

<~

0~~~--~,50----~,0-0~--~5n~~0

--

UJ

TEMPERATURE,

Fig. 20 The influence of prior plastic deformation ('JI) on martensitic transformation in an Fe--Cr+Ni alloy (17.2%Cr, 9.1%Ni). Deformation temperature ==+100°C, accelerating effect at low degree of deformation (up to 'I' = 8%) and retarding effect at higher 'P. (a) transformation during cooling to liquid nitrogen temperature and heating to room temperature; (b) isotherms at -125°C and -75°C for 'P = 8%, and 17%, (c) temperature dependence of the initial rate of isothermal A---7M transformation in undeformed (\{'= 0) and deformed alloy (8% and 17%103).

°C

20

o 401-------+--_+_---dr,~_t__--t__1

0:

~

~~~--4---~--~-¥--~

~

~-c

~-c ~

~20~--~--+---~~--t--1

-150

-100

tz

Vi 10 Annealed at 1000C

uJ

~200

~

i

uJ

-50

for 5h

--¥tt-'c----+----t

-\50 -\00 TEMPERATURE,

-c

-50

Fig. 21 Influence of neutron irradiation on A---7Mtransformation. Symbols denote different specimens tested. 't is time from irradiation to when cooling curves were obtained. (a) stimulating effect for a steel containing 0.48%C, 7.7%Mn, 2.2%Cu; (b) retardation effect for the Fe-Ni-Mn alloy (22.%Ni, 3.8%Mn). 't1 ==90 days, 't2 = 228 days, 't3 = 224 days.

142

Hatfield Memorial Lectures Vol. II

-IOr----~----r----...--. nvt=O

u

o

~

•. -40~~~-+--~~4---------~~

-60O~---2"'0-0---4"""O-O---6.&...OO---' ANNEALING

Fig. 22

TEMPERATURE,

°C

Displacement of the M, point of a Pe--Ni-Mn alloy (22%Ni, 3.%Mn), irradiated by neutrons, as a result of tempering at different temperatures.v!

distortions disappear only at high temperatures, and may be, for example, those on the boundaries of regions of coherent X-ray scattering, or some other kind of accumulated imperfection. These may hinder the cooperative displacements of atoms in nucleation or in the initial stages of growth, while the size of the martensite crystal is still too small to cause stresses sufficient to accelerate its further growth. Distortions of the austenite structure similar to those caused by plastic deformation arise as a result of partial transformation in the region surrounding a martensite crystal. That this facilitates the formation of new nuclei is illustrated in Fig. 23. It may be suggested that the well known phenomenon of austenite stabilisation is connected with a change in the distribution of these distortions and a decrease in the (associated) maximum stresses. A high density of structural imperfections arises in the austenite grain as a result of direct and reverse martensitic transformation. Their presence after reverse transformation can be easily shown by the marked increase in the response to etching of regions where austenite ~ martensite ~ austenite transformation has taken place (Fig. 24). The local stresses in this case cannot be very great because of the high temperature of the reverse transformation. Such structural changes in austenite grains make difficult both the formation of nuclei and their growth on subsequent cooling below M,62,63 (Fig. 25). The experimental data available show therefore that distortions of the regular crystal structure are very important in the kinetics of martensite transformation. The growth of martensite nuclei when their size becomes larger than the critical size can proceed more

Phenomena Occurring in the Quenching and Tempering of Steels

143

~ ~20~--~--~------~------r-----~

~ + 6 w

l-

v; Z

~ 15 tf---------c

a:=

L -.oJ

~ Q IO~--~--~------~------r-----~

TIME, min

Fig. 23 Increase in the rate of isothermal A~M transformation in an Pe--Ni+Mn alloy (22.4%Ni, 3.48%Mn) at T = -60°C after preliminary cooling to different temperatures T.l04

easily in a perfect austenite crystal. However, nucleation of martensite in such a crystal is extremely difficult, apparently because the formation of a nucleus by means of cooperative displacements in a perfect lattice requires a large energy due to the resistance of the elastic forces. Extreme supercooling below To is then necessary in order to permit the formation of martensite nuclei by thermal fluctuations. However, this supercooling may be so great that the mean thermal energy in this temperature range becomes too small for the necessary fluctuations to occur with sufficient frequency. For these reasons the formation of martensite nuclei in a perfect crystal of austenite will probably not take place at all, except perhaps in pure iron where To is high. Similar conditions may prevent the crystallisation of liquids. For example, after refining and de-activation of undissolved particles, salol does not crystallise, since the energy of nucleation becomes small enough in a temperature range where the atomic mobility is

144

Hatfield Memorial Lectures VoL II

Fig.24 Microstructure of an Fe-Ni-Co alloy (31.5%Ni, 5.5%Co) after direct and reverse martensite transformation. (a)relief arising from A~M transformation on cooling to -60°C; 20% of martensite was formed; (b) relief arising from the reverse M ~ A transformation. The same specimen as in (a) was repolished, heated to 580°C and cooled to room temperature; 0% of martensite; (c) the same as in (b) after repolishing and etching; the dark places are regions which have undergone A~M~A transformation.v? small.?" To cause pure salol to crystallise it is necessary to introduce solid particles able to facilitate nucleation, i.e. to decrease the energy of nuclei formation and hence the supercooling required. The temperature dependence of the rate of nucleation of martensite crystals indicates that for the formation of nuclei in 'previously prepared' places thermal fluctuations are

necessary. This can be shown especially clearly on steels with a low Ms. When transfor-

Phenomena Occurring in the Quenching and Tempering of Steels 145 60 ~,. A"--' /~

50

".tI'

(0) o

0--

~

~

40

Ihitial -.0

JJ'

o-. __

~

~'"

~Ip

x

~(o-~

0.."0..... Q,,"q.

w

X

~ 30

~a:::

w

-c ~ 20

V

j

10

o

~~

1'\,

...

...

~

L/. ~~ ~

-150

~

0

I

>I-t-J <1<1 ,"-J,

\

\

~

\.

M=30O/o M=45%

-50

o

125

Q

r=-;

\

(b)

150

0-

M=16%

or-

-100

°,..

175

100

75 50 25

4 .fl.";

Fig. 25 Stabilising action of preliminary A~M~A transformation on subsequent martensitic transformation in the Pe--Ni-Mn alloy (23.4%Ni, 3.3%Mn). M = quantity of martensite, obtained on preliminary cooling. (a) transformation of A~M on cooling to -194°C and on subsequent heating to +20°C; (b) temperature dependence of the initial rate of transformation. 105 mation is completely prevented by rapid cooling to liquid nitrogen temperature (Fig. 17), it implies that nuclei of critical size have no time to form in 'prepared' places while cooling, while at liquid nitrogen temperature the thermal energy is too small. Raising the temperature increases the probability of formation of the necessary fluctuation in these places and consequently the probability of nuclei formation. The rate of activation of 'prepared' places increases with temperature until the energy of nucleation begins to increase more rapidly than the mean thermal energy, and the transformation rate then starts to fall. Thus, in the formation of martensite nuclei in 'prepared places', the energy of atomic vibrations is, as in other phase transformations, the decisive factor. Structurally 'prepared' places, where nucleation is facilitated, play the same role here as undissolved solid particles in the crystallisation of liquids.

TEMPERING OF QUENCHED STEEL The phenomena of tempering in quenched steel are connected with the decomposition of the martensite, and, when a significant quantity of retained austenite is present, with the decomposition of this constituent also. The supersaturated solid solution of carbon in a-iron is highly unstable so that its decomposition starts even at room temperatures, and the structural changes connected with its transition to a two-phase mixture of a-iron and

146

Hatfield Memorial Lectures VoL II

cementite, proceed at all temperatures from room temperature to the Ai point. To every tempering temperature there corresponds a certain state of the steel, which is reached in a comparatively short period of time, and which changes only very slowly on further tempering. This state usually changes gradually with increasing tempering temperatures, but there are comparatively narrow temperature intervals in which the change is considerable. These intervals are marked by effects on the curve showing the variation of different properties on heating. The processes taking place in these temperature intervals are called the first, second and third transformations on ternpering.v'' The position of these intervals depends upon the heating rate; at a heating rate of about 10 K min-1 they lie at about 100-lS0°C, 2S0-32SoC and 32S-400°C. These transformations can be practically completed within a few hours at temperatures of 110, 2S0 and 325°C, but at lower temperatures they require an extremely long period of tempering. Apart from the effects at these temperatures another distinct change in mechanical properties is observed at about 4S0°C.66 The changes in the properties of quenched steel on- tempering are mostly due on the one hand, to changes in the state of the a-solid solution and, on the other, to the formation and changes of state of the carbon rich phase; however, the effects of 'the second transformation on tempering' are governed by the decomposition of the retained austenite.

Changes of State of the a-solid Solution The Two Stages of the Decomposition of Martensite X-ray investigations of the changes in the crystal structure of quenched steel associated with changes in properties at tempering temperatures of 100-lS0°C show that the latter are connected with a decrease in the axial ratio of the tetragonal lattice of martensite, and consequently with precipitation of carbon from solid solution.4,6,7,67,68 As a result of this process a non-uniform tetragonal structure with a small axial ratio is formed. 67,5 This process is the first stage of martensite decomposition. Tempering at temperatures of 150300°C leads to a gradual decrease of non-uniformity and axial ratio, i.e. to further gradual precipitation of carbon from solid solution. These general conclusions follow from the study of line broadening in tempered steels with carbon contents ranging from 0.1 to

1.4%.69 Both stages of martensite decomposition have been clarified by investigating the 9 tempering of quenched single crystals. ,13 Owing to the regular orientation of the martensite lattice with respect to the initial austenite it is possible to avoid difficulties associated with the overlapping doublet components due to tetragonality. It appears that the first stage of decomposition has a so called 'two-phase' character. Apart from the solid solution of initial concentration an a-solid solution, containing about 0.3%C, appears and during tempering the volume of this second solid solution increases at the expense of the initial one.l ' The kinetics of the 'two-phase' decomposition have been investigated at temperatures of 80, 100 and 120°C.70,71 The rate of the process can be characterised by

Phenomena Occurring in the Quenching and Tempering of Steels 147 the time of half decomposition, that is, the time when the volume of both solid solutions becomes equal. Table 3 shows the results of determinations of this time at different temperatures for a 1%C steel. For the above mentioned temperatures and for room temperature the data are obtained directly from experiment; the other values are calculated from the temperature dependence of the rate of decomposition. Table 3 t,OC

o

20 40 60

80

Time of half decomposition of martensite in the first stage of decomposition 'to.s

t,OC

'to.s

340 years 6.4 years 2.5 months 3 days 7 h 50 min

100 120 140 160

50 min 8 min 2.3 min 45 s

71

While investigating martensite powder which had been isolated electrolytically, it was found that in addition to two-phase decomposition there is also one-phase decomposition.72 The former was established from the change in relative intensity of the lines of tetragonal doublets, and the latter from the change in their separation (similar to that observed in tempering quenched steel powder). 7 Thus the character of the first stage of decomposition may depend on the conditions. The solid solution arising as a result of the first stage has a tetragonal lattice with an axial ratio 1.012-1.013. This axial ratio corresponds to about 0.2S-0.3%C.13 The solid solution with this concentration possesses a stability considerably higher than the initial one, and further decrease of carbon content in solution at temperatures of 100-lS0°C proceeds extremely slowly.T' The process of further gradual carbon precipitation representing the second stage of the martensite decomposition extends to temperatures of lS0-300°C. Within the first few hours tempering, the carbon content reaches a certain value which changes only slightly thereafter.69,73 The quantity of carbon remaining in solution for different tempering temperatures has been estimated fromthe changes in the breadth oflines,69,73 the specific column.?? and the axial rario.l ' The precision of these estimates is not high; the values obtained from the line breadth are somewhat exaggerated and those obtained from the axial ratio are too low (Table 4). After 1 h tempering at 300°C, about O.l%C apparently still remains in solid solution in medium and high carbon steels. In steels containing 0.4%C and less, the first transformations on tempering (100-1S0°C) have no effect on the curves showing the variation of properties. The absence of a decomposition process is confirmed by the curves showing the variation of the line breadth on tempering (Fig. 26). Obviously in these steels the two-phase stage of decomposition is absent and a gradual precipitation of carbon takes place which extends over a wide range of temperature. The data on the changes of line breadth on tempering show that in quenched steel containing 0.1%C precipitation of carbon from solid solution starts only at temperatures above 250°C, and in steels having 0.2% and 0.3%C, above 200°C and 1S0°C respectively.

148

Hatfield Memorial Lectures VoL II

Table 4 Concentration of carbon retained in a-solid solution after tempering 1 h at different temperatures in the second stage of martensite decomposition. Values obtained by various methods Tempering temp.,oC

Tempering time

Axial ratio13

RT 100 125 150 175 200 225 250

10 years 1h

0.27, 1.4 0.29,1.2 0.29 0.27 0.21 0.14 0.08 0.06

Carbon content, 0/0 Specific volume?"

Line breadth69

0.42 0.37

0.52 0.45 0.39 0.32 0.28

0.32 0.15

4'0

E

E

w

~ 3'0 - ---~-"'oR..---\-"'II--------+---

u..

--------'f-------i

o ::c

o -c w

a::

C(l

Fig. 26

Change

in breadth

of the line (112), (Cr radiation) different carbon contents.v?

on tempering

steels with

Influence of alloying elements on the process of the precipitation of carbonfrom a-solid solution The rate of the first martensite decomposition changes little in the presence of alloying elements (chromium, molybdenum, tungsten, silicon), although the half decomposition time somewhat increases. Indeed, to obtain the same time of half decomposition as in carbon steel, it is necessary to change the tempering temperature only by a few degrees.f" On the other hand, the presence of alloying elements can lead to a great increase in the temperature range of the second stage of decomposition, e.g. by as much as 100 or 200°C.

Phenomena Occurring in the Quenching and Tempering of Steels 149 only below 275molybdenum, tungsten, titanium, vanadium, silicon, cobalt) the tetragonal lattice can be revealed after tempering at 400, 450 and even up to 500°C (Table 5, Fig. 27).14,87,88 From this it follows that the presence of these elements leads to an increase in the stability of the supersaturated solid solution of low carbon content. On the other hand, manganese and nickel somewhat decrease its stability. This conclusion has been confirmed by data on the tempering of alloy steels containing less than 0.2%C (Fig. 28). In steels with O.l%C, containing titanium, vanadium and molybdenum, decomposition is retarded up to 400SOO°C,15 and in the range SOO-SSO°C decomposition proceeds with the formation of special carbides and with rising hardness (secondary hardening (Fig. 29)). The high stability of low carbon steels (0.1-0.2%C) containing certain alloying elements was also used for studying the mechanism of the isothermal decomposition of austenite in the intermediate (bainitic) region.89,90 For carbon steel the tetragonal lattice can be observed at temperatures

300°C. For steels containing some alloying elements (such as chromium,

Microstresses and coherent regions The use of single crystals and the selection of reflections with the third index equal to zero make it possible to avoid the influence on line broadening of the overlapping of the doublet lines of the tetragonal lattice, and the non-uniformity of concentration of the solid solution.F! Investigation showed that these two factors have practically no influence '·14~----,.------,-----.,....------' o 1-4OJo C x ~05% C, /-38 n %

-'.Q9%C,I-70%Mn

/'/2~-----"~~-----t-

-o

u

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Fig. 27

Influence ofTi and Mn on the course of the second stage of the decomposition

martensite.

14

of

150

Hatfield Memorial Lectures VoL II 0~~C\J • ex> ('t)

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Phenomena Occurring in the Quenching and Tempering of Steels

151

50

I

u

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'\

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200 400 600 ANNEALING TEMPERATURE.oC

of Ti on the decomposition of martensite O.1%C.1S

on tempering

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Tempering

200

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400

TEMPERATURE,O(

600

of quenched iron containing O.04%C.

on the broadening of the lines of polycrystalline steel at temperatures higher than 250°C. The effect of tempering on both the microstresses and the size of the regions of coherent X-ray scattering in carbon steels, were estimated approximately by using the different angular dependence of the line broadening caused by these two factors (Fig. 30). The size of coherent regions starts to increase only at temperatures higher than 300-3S0°C, and they still cause some line broadening even on tempering at SOO-SSO°C. Microstresses decrease somewhat at lS0-200°C, but they still remain high up to 300°C. In the interval 300-4S0°C the microstresses decrease rapidly, and after tempering at 500°C they are not observable by the method used. On tempering in the range 3S0-400°C, the line broadening approaches that which arises from the cold plastic deformation of anriealed stee1.91,92 In steels with alloying elements delaying the precipitation of carbon from solid

solution, the fall in the stresses is displaced towards higher temperatures.

93

152

Hatfield Memorial Lectures VoL II

o

Fig. 30

Change of microstrains l1al a and of size of coherent regions (D) on tempering 1.4%C steel. 21

a

Formation of Carbide Phase on Tempering As in easily determined by X-ray diffraction="? the carbide phase formed as a result of tempering above 300°C is the carbide Fe3 C (cementite). The investigation of the ternpering of single crystals enabled the orientation of the cementite lattice with respect to those of the austenite and martensite to be determined. 75, 76 The relation can be expressed as follows:

(103) Fe3 C [010] Fe3 C

II II

(Oll)M [lll]M

II II

(lll)A [110]A

This gives 12 positions of the cementite lattice with respect to that of austenite. By comparing X-ray photographs of carbide powders extracted electrolytically from a quenched high carbon steel, tempered at various temperatures from 200 to 700°C, it was found that on gradually decreasing the tempering temperature below 600°C, the line breadth of cementite increases. Also the lines with a high third index become much broader than the others, and as the tempering temperature decreases they appear gradually weaker, ultimately merging with the diffuse background and becoming invisible.Tr"? (Some strong lines with similar angles of reflection merge together here.) The changes in the X-ray photograph are shown schematically in Fig. 31.

Phenomena Occurring in the Quenching and Tempering of Steels 153

. I

II

Fig. 31

bOO°C

Relative line intensity on X-ray photographs of cementite powder extracted from quenched steels, tempered in the temperature range 200-700°C.79

The strong broadening of lines with a high third index indicates that the cementite crystals are small in the direction of the c axis and so have a plate-like form. By measuring the line breadth the dimensions of the cementite crystals along the three axes were estimated (Fig. 32). At tempering temperatures from 200 to 350°C the size of the particles along the c axis is about 10 lattice spacings, and along the perpendicular directions about 40. At temperatures higher than 350-400°C growth of the crystals is observed. It is to be noted that alloying elements which cause the second stage of martensite decomposition to be delayed also displace up to 450-550°C the temperatures at which cementite crystals start to grow (Fig. 33).79 X-ray study of the carbide phase in bulk specimens of quenched steel tempered at low temperatures (lower than 300°C) encounters considerable difficulties, but these have been overcome to a certain extent by using austenite single crystals,80,81 monochromatic

radiation.s- and by means of electron diffraction.83-85

In the first method features of the

154

Hatfield Memorial Lectures VoL II 10'01~----~------~----~------~

b· E

u -D

Q

400

500

bOO

TEMPERATURE,oC

Fig. 32 Change of dimensions of cementite particles with increasing tempering temperature.?? mxA, m)3, mzC are size of cementite crystals in the directions A, Band C; A, Band C are lattice parameters of cementite. diffraction pattern of the carbide phase were observed which differed from those of Fej C. This was considered to confirm the existence at low tempering temperatures of a carbide different from cementite and which was called FexC.80 The features due to FexC disappear on tempering in a temperature interval of300-350°C. Photographs after low temperature tempering both using monochromatic X-rays and also using electron diffraction showed lines of the carbide phase at the same angles of reflection, as on the photographs of single crystals. The observed lines may be related to a hexagonal lattice similar to that ofFe3N.82-85 These data suggest that on low temperature tempering, as a result of the first stage of martensite decomposition, an intermediate E -carbide with a hexagonal lattice precipitates.

Nature of the Phenomenon of Tempering On tempering decomposition

quenched steel the main process is the decomposition of martensite, the of a supersaturated solid solution of carbon in a-iron. Therefore there

Phenomena Occurring in the Quenching and Tempering of Steels 10~~----~------~----~----~----~ /

200 Fig. 33

300

400

500

TEMPERATURE;t

155

/

bOO

700

Influence of alloying elements on the growth of cementite particles on tempering. 79

must occur on tempering the same process that takes place on ageing supersaturated solid solutions where the solubility of the second component increases with rising temperature. In general this is the process of the precipitation of a dissolved element and formation of a second phase under conditions where the diffusion rate is slow. However, although the martensite decomposition is in many ways similar to the decomposition of other supersaturated solid solutions, it possesses its own peculiarities which strongly distinguish it from others. The reasons for these peculiarities are as follows. First, the supersaturated solid solutions is not obtained here by a simple 'freezing' of a high temperature state, but is a result of a diffusionless transformation. This leads to a fine micro- and submicrostructure and to the presence of different kinds of inhomogeneities and imperfections. Second, the mobility of the atoms of the solvent and the dissolved elements differ considerably from each other. Third, the properties of the crystals of the solid solution depend greatly on the concentration of the element dissolved. The high supersaturation of carbon in the a-phase in medium and high carbon steels causes it to be highly unstable and leads to a first stage of decomposition at 10O-150°C.

156

Hatfield Memorial Lectures Vol. II

The structure fonned as a result of the first stage of martensite decomposition is called tempered martensite. Steel in this state possesses almost the same high hardness as in the quenched state, but, however, has higher ductility. X-ray investigations show that tempered martensite itself consists of a partly decomposed a-solid solution. The martensite crystals still contain a considerable amount of carbon in solution, and dispersed carbide particles formed as a result of decomposition are uniformly distributed inside them. 67,69,13 The state of the tempered martensite gradually changes at tempering temperatures of 150-300°C and the carbon content in solution decreases (Table 3). Particles of cementite appear, and may be observed by means of the X-ray diffraction pattern in a steel tempered at 200°C; their quantity increases on raising the tempering temperature to 300°C. The concentration remaining in solid solution depends only slightly on the carbon content for medium and high carbon steels, but it is the greater, the greater the initial concentration. The difference in amount of the dissolved carbon for various steels decreases with increasing temperature (Fig. 26). This conclusion is confirmed by specific heat data (Fig. 34i) obtained on heating quenched steels.?" The austenite decomposition causes a large thermal effect above 250°C, which is superimposed on the effect due to the second stage of the martensite decomposition. If, before obtaining a curve, the quenched steel is first tempered at 250°C, the thermal effects due to the first and second transformations disappear, (i.e. effects due to the first stage of the martensite decomposition, and of the decomposition of the retained austenite (Fig. 34ii)). These curves show an effect at 250-325 °C and, well separated from it, an effect at 325-400°C. Similar curves are obtained on heating quenched steels containing 0.4% of carbon and less (Fig. 34i). The first effect is apparently governed by precipitation of the dissolved carbon, and its magnitude, for the steels given a preliminary temper (Fig. 34ii), is nearly independent of the carbon content of the steel and is about that for a quenched steel with 0.22%C. Thus even after tempering at 250°C there still remains about 0.2%C in the a-solid solution. Crystals of martensite with such a carbon content in solid solution themselves possess a high elastic limit; this can be further increased considerably, owing to the presence of dispersed carbide particles inside them. Therefore, after such tempering, the hardness of steel still remains high. As has already been stated, in quenched low carbon steels the first stage of martensite decomposition is absent. However, during tempering in this range of temperature some properties of these steels can change considerably; this may be connected with a different kind of relaxation process. For example, after tempering quenched steel with 0.1 %C at 200°C the coercive force decreases by as much as twofold, while the hardness, line breadth and specific volume start to change at considerably higher temperatures. One can guess that the decrease in the coercive force is connected with the relaxation of microstresses caused by diffusion from regions of compression to regions of tension.95-97 An elastic deformation becomes in part a non-elastic one. A considerable increase in the strength of quenched low carbon steels with 0.1-0.2% which takes place on holding at low temperatures (beginning at room temperature) 98 is apparently also connected with the relaxation of stresses.

Phenomena Occurring in the Quenching and Tempering of Steels

157

(i)

Fig. 34

Specific heat curves obtained in heating steels with different carbon contents. (i) after quenching; (ii) after quenching and tempering for 2 h at 250°C.

The mechanism of decomposition of the retained austenite on tempering is the same as that of the isothermal decomposition of austenite in the intermediate temperature range. The products of decomposition are similar to those of martensite tempered at the. same temperatures. A rapid decrease of hardness starts on tempering above 300°C, and on tempering between 300 and 400°C considerable changes take place in some of the properties of steel. This change of state is called the 'third transformation' on tempering. In this temperature range, it is to be noted that large changes in specific volume of steel occur without noticeable changes in the lattice constant of the a-phase. Figure 35 shows data?" on the change of specific volume with the carbon content for the first (~V1) and third (~V3) transformations, and also on the total change (~V:) in the specific volume of quenched steel on tempering to 500°C. The volume change due to the third .transformation itself represents a considerable part of the whole effect on tempering. Figure 36

shows how the thermal effect of the third transformation depends

on the carbon

158

Hatfield Memorial Lectures Vol. II

content.?? The disappearance, just in the temperature range of 350-400°C,80,81 of particular lines of the carbide formed during low temperature tempering suggested that the effects of the third transformation are to be attributed to the transition of a low temperature carbide to cementite. However, further investigations?" showed that the cementite is already present after tempering at 200°C, and exists in considerable quantity before the beginning of the 'third transformation'. Apparently at tempering temperatures of 200300°C both the E -carbide and cementite are present. Measurements of the lattice constants of cementite powders extracted from steel tempered below 300°C show that the volume of the elementary cell of this cementite is less by 3.8% than that of normal cementite.U'? It is possible that this cementite contains less carbon than normal cementite. Also the constants of the cementite lattice in a specimen tempered below 300°C may be changed due to coherence with the a-phase. At present there are no experimental data about the structural changes on tempering in the range of 300-400°C. To come to definite conclusions on the nature of the third transformation it is necessary to determine whether the volume changes accompanying it are affected by microscopic cracks and pores arising on quenching, and also in particular, on low temperature tempering. In alloy steels a considerable upward displacement of the temperature ranges of tempering processes can occur. This applies especially to the second stage of the first transformation on tempering; to the beginning of growth of the carbide particles; and to the third

~ 100

---------+---~~-+-----_+_-__i

0·4

C,

0-8

1'2

%

Fig. 35 Change in specific volume of quenched steels on tempering as a function of carbon content. ~ is total volume change on tempering; ~ V1 effect of 'first transformation'; ~ V3 is effect of 'third transformation'.

v:

Phenomena Occurring in the Quenching and Tempering of Steels 159 (·4 ~J'2

"8

1·0

tj' w

0'8

~ 0·6 ~ 0·4

Fig. 36

~

0'2

~

0

Dependence

~.

~

0·2

-~

-' ~

~ ~

0-4

0·6

c,

o.e

-~-

1-0

»>

1-2

j-4

%

of the thermal effect of the 'third transformation' content.

on the carbon

transformation on tempering. The state of the tempered martensite in these steels, and the high hardness, are retained to higher temperatures. In tempering at low and medium temperatures there is almost no redistribution of alloying elements during tempering because of the small mobility of metal atoms below 400-S00°C. The concentration of alloying elements in the carbide phase is the same as in the a-solid solution, and the slow rate of diffusion of the alloy elements at low temperatures causes one of the intermediate carbide phases to be an alloyed cementite with the same concentration of alloying elements as in martensite. The more stable special carbides are formed only on tempering at temperatures higher than SOO-SSO°C. At low carbon contents in some alloy steels (less than O.2%C), the absence of diffusion of alloying elements is obviously responsible for the fact that the decomposition of the martensite and the decrease of hardness, do not take place up to 4S0-S00°C. We may suppose that in these steels the formation of a heterogeneous state, a+ alloyed cementite requiring only the diffusion of carbon is not stimulated thermodynamically. For the formation of the state, a+ special carbide, which is thermodynamically more stable compared with the initial martensite, it is necessary to have diffusion of alloying elements. This process proceeds at temperatures higher than 500°C causing the phenomenon of 'secondary' hardening. is

REFERENCES 1. 2. 3. 4.

A. SAUVEUR: Trans. AIME, 1926,73,859-908. W. L. FINK and E. D. CAMPBELL: Trans. Am. Soc. Steel Trat., 1926,9,717-754. N. SELjAKOV et al.: Zhurnal prikladnoijiziki, 1927, (2),51; Z. Physik, 1927,45,384-408. G. KURDjUMOV and E. KAMINSKIY: Zhurnal prikladnoijiziki, 1929, (2),47; Z. Physik, 1929, 53,696-707. 5. A. IVENSEN and G. KURDjOMUV: Zhur Fiz. Khim., 1930, (1),41. 6. E. OHMAN:]ISI, 1931, 123,445-463. 7. G. HAGG:]ISI, 1934,130,439-451. J

160

Hatfield Memorial Lectures VoL II

8. G. V. KURDjUMOV: Sbornik dokladov sektsii metallovedeniya i termicheseoi obrabotki, VNITO metallurgov, Moscow, 1940,96. 9. G. KURDjUMOV and G. SACHS:Z. Physik, 1930,64,325-373. 10. G. V. KURDjUMOV and N. L. OLSON: Zhur. Tekhl1. Fiz., 1939,9,1891. 11. G. V. KURDjUMOV et al.: Stal', 1935, (4), 84. 12. E. S. KAMINSKYand M. D. PERKAS:Problemy metallovedeniya ifiziki nietallov, 1949, (1),211. 13. G. V. KURDjUMOVand L. 1. LYSAK:Zhur. Tekhn. Fiz., 1946, 16, 1307;J1S1, 1947, 156, 2936. 14. G. V. KURDjUMOV and M. D. PERKAS:Problemy metallovedeniya ifiziki metallov, 1951, (2), 153. 15. M. D. PERKAS:Problemy metallovedeniya ijiziki metallov, 1952, (3), 139. 16. R. 1. ENTIN: Problemy metallovedeniva ifiziki metallou, 1949, (1),281. 17. L. I. KOGAN and R. I. ENTIN: Problemy metallovedeniya ifiziki metallov, 1951, (2), 216. 18. G. V. KURDjUMOV: Problemy metallovedeniya ifiziki metallov, 1952, (3),31. 19. P. L. GRUSIN et al.: Doklady AN, 1953,93, 1021. 20. G. V. KURDjUMOV and M. D. PERKAS:Doklady AN, 1956, 3, 818. 21. G. V. KURDjUMOV and L. I. LYSAK:Zhur. Tekhn. Fiz., 1947, (17),993. 22. M. P. ARBusov: Doklady AN, 1950,74,1085. 23. M. P. ARBusov et al.: Doklady AN, 1953, (90),375. 24. M. P. ARBusov: Voprosy Fiziki metallov i metallovedenya, 1955, (6),3. 25. G. V. KURDjUMOV: Zhur. Tekhn. Fiz., 1954, (24), 1254; Problel1ty metallovedeniya i fiziki metallov, 1955 (4),321. 26. V. M. GOLUBKOVet al.: Fiz. Met., 1957, (5), 465; Problemy metallovedeniya ifiziki metallov, 1958, (5), 433. 27. L. S. MOROS: Zhur. Tekhn. Fiz., 1952, (22),498. 28. H. LIPSON and A. M. B. PARKER:J1S1, 1944, 149, 123-141. 29. V. A. ILJINAet al.: Doklady AN, 1952, (85), 197. 30. V. K. KRIZKAjAet al.: Zhur. Tekhn. Fiz., 1955, (25), 177. 31. V. K. KRIZKAjAet al.: Fiz. Met., 1958, (6), 177. 32. G. V. KURDjUMOV et al.: Fiz. Met., 1959, (7),747. 33. V. M. KARDONSKYet al.: Fiz. Met., 1959, (7),752. 34. S. S. STEINBERG:Meta llurg, 1937, (10), 58. 35. G. V. KURDjUMOV: Zhur. Tekhn. Fiz., 1948, 18, 999; Problemy metallovedeniya i fiziki metallov, 1949, (1), 132. 36. A. R. TROIANO and A. B. GRENINGER: Met. Prog., 1946,50,303-307. 37. B. A. BILBYand]. W. CHRISTIAN: lnst. Metals Monograph and Report Series no. 18, 1955, 121. 38. G. V. KURDjUMOV: Isvest. Akad. Nauk., Fiz. mat., 1936, (2),271. 39. C. ZENER: Met. Tech., 1946,13, pt.1, 1-22. 40. L. KAUFMANand M. COHEN:]. Met., 1956,8, 1393-1401. 41. G. WASSERMANN:Metallwirtschaft, 1934,13,133-140. 42. E. KAMINSKYet al.: Zhur. Tekhn. Fiz., 1934, (4), 1774. MetallwirtschaJt, 1934, 13,373. 43. V. GRIDNEV and G. KURDjUMOV: Zhur. Tekhn. Fiz., 1937, (7),2091; Tekll1l. Fiz. USSR, 1938,5, (1). 44. V. GRIDNEV: Metallurg, 1938, (4).

Phenomena

Occurring in the Quenching and Tempering of Steels

161

45. V. GRIDNEV: Zhur. Tekhn. Fiz., 1941, (11), 1226. 46. jA. M. GOLOVCHINER: Problemy metallovedeniya ifiziki metallov, 1951, (2), 119. 47. G. V. KURDjUMOV and L. G. KHANDROS: Zhur. Tekhn. Fiz., 1949, (19),761. 48. E. SCHElL: Z. anorganische Chemie, 1929, 183,98; 1932,207,21. 49. J. B. HESS and C. S. BARRETT:]. Met., 1952,4, (6),645-650. 50. I. V. ISAICHEV et al.: Trans. AIME, 1938, 128,361-367. 51. G. KURDjUMOV et al.: Zhur. Tekhn. Fiz., 1938,8, 1959; Zhur. Fiz., USSR 1940, 3, 297308. 52. G. KURDJUMOV and E. KAMIHSKY: Zhur. Tekhn. Fiz., 1936, (6), 987; Metallwirtschqft, 1936, 15,905. 53. G. KURDjUMOV: Trans. AIME, 1939, 133,222-223. 54. G. V. KURDjUMOV and O. P. MAKSIMOVA: Doklady AN, 1948, (61),83. 55. jA. M. GOLOVCHINER and G. V. KURDjUMOV: Problemy metallovedeniya i fiziki metallov, 1951, (2),98. 56. G. V. KURDjUMOV and O. P. MAKSIMOVA: Doklady AN, 1951, (81),565. 57. G. V. KURDjUMOV and O. P. MAKSIMOVA: Isvest. Akad. Nauk, Otdel. Tekhn., 1957, (6),4. 58. G. V. KURDjUMOV: Present-day metallurgical problems (Sovremennye problemy metallurgii), 34; Akademizdat 1958, Moscow;]. Met., 1959, july, 449-453. 59. B.JA. LjUBOV and A. L. ROITBURD: Doklady AN, 1958, (120), 1011.. · 60. G. V. KURDjUMOV et al.: Doklady AN, 1957, (114), 768. 61. A. I. SAKHAROV and O. P. MAKSIMOVA: Izvest. Akad., Nauk. Otdel. Tekhn., 1958, (7),3. 62. O. P. MAKSIMOVA and A. I. NIKONOROVA: Problemy metallovedeniya ifiziki metallov., 1955, (4), 123. 63. JA. M. GOLOVCHINER andJu. D. TjAPKIN: Doklady AN, 1953, (93),39. 64. V. I. DANILOV: Problemy metallovedeniya ijiziki metallov 1949, (1),7. 65. H. HANEMAN and L. TRAGER: Stahl Eisen, 1926,46, (2), 1508-1514. 66. K. F. STARODUBOV: DokladyAN, 1946, (53),217. 67. G. V. KURDjUMOV: Zhur. Fiz. Khim., 1930, (1),281; Zhur. Fiz., 1929~55, 187-198. 68. G. V. KURDjUMOV: Vestnik metallopromyshlennosti, 1932, (9), Arch. Eisenhiitt., 1932-33, 6, 117-123. 69. G. V. KURDjUMOV and N. L. OSLON: Zhur. Tekhn. Fiz., 1939, (9), 1891. 70. G. V. KURDjUMOV and L. 1. LYSAK: Zhur. Tekhn. Fiz., 1949, (19),525. 71. L. I. LYSAK: Voprosy fiziki metallov i metallovedeniya, 1952, (3), 46. 72. M. P. As.aosov: Voprosy jiziki metallov i metallovedeniya, 1955, (6), 3. 73. E. S. KAMINSKY and T. I. STELLEZKAjA: Problemy metallovedeniya ijiziki metallov, 1949, (1), 192. 74. E. S. KAMINSKY and D. KAZNELjSON: Zhur. Tekhn. Fiz., 1945, (15), 182. 75. M. P. ARBUSOV and G. V. KURDjUMOV: Zhur. Tekhn. Fiz., 1941, (11),412; Zhur. Fiz. USSR, 1941,5. 76. I. V. ISAICHEV: Zhur. Tekhn. Fiz., 1947, (17),835. 77. M. P. An.nusov: Zhur. Tekhn. Fiz., 1949, (19), 1119. 78. M. P. An.nusov: Doklady AN, 1950, (73),83. 79. M. P. Arususov: Voprosy jiziki metallov i metallovedeniya, 1952, (3),3. 80. M. P. ARBUSOV and G. V. KURDjUMOV: Zhur. Tekhn. Fiz., 1940, (10), 1093; Zhur. Fiz., 1941, 5, 1093.

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81. 1. V. ISAICHEV:Zhur. Tekhn. Fiz., (17), 1947, 839. 82. K. H.JAcK:JISI, 1951,169,26-36. 83. K. D. HEIDENREICH et al.:J. Appl. Phys., 1946, 17,127-136. 84. J. TROTTER and D. McLEAN:jISI, 1949, 163,9-13. 85. Ju. A. SKAKOVet al.: Doklady AN, 1958, (118), 284. 86. G.JA. KOSYRSKYand G. V. KURDjUMOV: Voprosyjiziki metallov i metallovedeniya, 1950, (2), 38. 87. 1. V. ISAICHEV and E. S. KAMINSKY: Trudy Instituta Metalhl1~ii AN USSR, 1946; G. V. KURDjUMOV: Informatsionnyi byulleten, AN Ukr. SSR, 1943, (5), 13. 88. L. 1. LYSAKand G. JA. KOSYRSKY: Voprosy jiziki metallov i metallovedeniya, 1952, (3),53. 89. L. 1. KOGAN and P. I. ENTIN: Problemy metallovedeniya ifiziki metallov, 1958, (5), 161. 90. G. V. KURDjUMOV and M. D. PERKAS: Problemy metallouedeniya ijiziki metallov, 1951, (2), 167. 91. E. S. KAMINSKYetal.: Zhur. Tekhn. Fiz., 1941, (11), 1089. 92. G. V. KURDjUMOV: Voprosy jiziki metallov i metallovedeniva, 1950, (2),3. 93. L. I. LYSAKand E. G. NESTERENKO: Voprosy fiziki metallov i metallovedeniya, 1953, (4), 12. 94. P. L. GRUSIN et al.: Metallurg, 1940, (8), 15. 95. I. A. BILjDSjUKEVICHet al.: Problemy metallovedeniya ifiziki metallov, 1955, (4), 205. 96. N. S. FASTOV: Problemy metallovedeniya i jiziki metallov, 1955, (4), 219. 97. JA. M. GOLOVCHINER and V. M. GOLUBKOV: Problemy metallovedeniya i fiziki tnetallov, 1955, (4), 222. 98. V. I. SARRAK and R. I. ENTIN: Doklady AN, 1959, (127),306. 99. R. I. ENTIN: Problemy metallovedeniya ijiziki metallov, 1955, (4),239. 100. G. V. KURDjUMOV and R. I. ENTIN: Otpushchnaya khrupnost' konstruktsionnykh stalei, 1945, Metallurgizdat.

101. G. V. KURDjUMOV et al.: Problcmv metallovedeniya ifiziki metallov, 1955, (4),228. 102. G. V. KURDJUMOV and L. G. KHANDROS: Doklady AN, 1949, (66),211. 103. G. V. KURDjUMOV et al.: Fiz. Met., 1958, (6),95. 104. O. P. MAKSIMOVAand E. O. ESTRIN: Fiz. Met., 1960, (9),426. 105. I. M. SAHER et al.: Izvest. Akad. Nauk. Otdel. Tekhn.: Metallurgya i toplivo, 1960, (2). 106. G. V. KURDJUMOV et al.: Doklady AN, 1950, (73),307. 107. N.J. PETCH:jISI, 1943,147,221-227. 108. F. E. WERNER: et al.: Trans. ASM, 1957,49,823-841. 109. M. P. AZBUSOV: Private communication.

THE

FIFTEENTH

HATFIELD

MEMORIAL

LECTURE

Metallography - A Hundred Years after Sorby A. G. Quarrell At the time the lecture was given, Professor Quarrell was Professor of Metallurgy at Shrffield University and Dean of the Faculty of Metallurgy, also at Sheffield University. The lecture was presented at Firth Hall, University of Shrffield, on the evening of 8 May 1963. The prologue to Professor Quarrell' s lectureforms the Introduction to the present volume.

Born on the outskirts of Sheffield in 1826, Henry Clifton Sorby came from a long line of cutlers. Thanks to the family business he was of independent means and so was able to devote his life to the scientific studies to which he had been introduced by his tutor, the Reverend Walter Mitchell. Sorby was a brilliant amateur who .made original contributions to archaeology, biology, chemistry, geology and meteorology, as well as to metallurgy. In all 'these fields his work was noteworthy because of the brilliance of his techniques, but he was interested in techniques chiefly because of the information they could provide. He was quick to see, and to follow up, potential applications of his techniques. This took him into diverse fields which included the detection of blood, sewage contamination and the purifying action of minute animals and plants, and the causes of colour in such varied things as autumn leaves, clouds and sky, human hair, algae and birds' eggs. Of no man could it more truly be said that he was concerned with the nature of things.

w. C.

Williamson, a Manchester surgeon, had developed a technique for preparing thin sections of biological materials for microscopic 'examination. He taught it to Sorby, who in 1849 extended it to the Geological Society in 1850. Primarily because of his interest in meteorites and-his desire to explain their structure, Sorby adapted his techniques to the examination of the iron and steel products of 'his native city. He was not the first to examine a metal surface 'under a microscope.P even a polished and etched metal surface," but he was the first to make a systematic study of the way in which structure varied with composition, heat treatment and manufacturing process, and to appreciate that incorrect surface preparation could give misleading results. He is generally recognised as the father of metallography. _ Sorby gave an account of his work on metals to the Sheffield Literary and Philosophical Society in the spring of 1864, and to the British Association Meeting in the autumn of the

Or

163

164

Hatfield Memorial Lectures VoL II

Dr Henry Clifton Sorby.

Metallography - A Hundred Years after Sorby

165

same year. He fully described his techniques in a chapter of Beale's book on microscopy'' published in 1868. In cooperation with a local photographer, Charles Hoole, he photographed some of his microstructures at a magnification of nine diameters. These were shown at the British Association Meeting and later presented to the Science Museum (1876). Y et for several years English metallurgists showed little interest in his work, and Sorby turned his attention to marine zoology. Because of the later metallographic work of.Martens= and of Wedding? Sorby's interest in the subject was stimulated once more, and in 1882 he was again lecturing on the work he had done nearly 20 years earlier. He also accepted an invitation to describe his work at the 1885 Annual General Meeting of the Iron and Steel Institute, and this was followed by two detailed papers in the Journal in 1886 and 1887.9,10 In spite of the limited period during which he worked on iron and steel, Sorby learnt much about them, and his publications show a surprising understanding of the phenomena he encountered. Y et several years later metallography was in such a poor state of development in Sheffield that Arnold"! considered it necessary to devote an evening meeting of the Sheffield Society of Engineers and Metallurgists to a detailed description of the preparation of microspecimens, and to appeal for recruits to learn it. Even he referred to it as 'an apparently unpractical subject', though he expressed his conviction that it was important to the future of the steel industry. Unlike some of the workers who followed him, Sorby was not satisfied merely to give a careful descripton of his metallographic observations; it is typical of him that he tried to explain them, and considered their implications. He not only recognized six 'well defined constituents' in steels, but realised how these might be responsible for the observed properties. By observing the way in which the microstructure changed after various treatments, he came to understand the nature of decarburisation, of graphitisation and of recrystallisation in welding, in cooling from high temperatures and in annealing after cold work. He recognised the cold worked state as one of unstable equilibrium, as the following quotation shows: . . . when distorted the particles must be in a state of unstable equilibrium and we can therefore readily understand why recrystallisation so easily takes place whenever the circumstances are such as to permit the particles to rearrange themselves in a state of stable equilibrium. It appears to me that this is a general principle of great importance in connexion with the mechanical properties of worked iron, probably often overlooked. On the fatigue of metals he wrote in his 1887 paper: It was at one time supposed that by continual vibration a bar of so called fibrous, iron becomes crystalline. To test this question, a bar was fixed on a tilt hammer in such a way as to vibrate up and down continuously for 15 hours until it broke .with a crystalline fracture. A longitudinal section of the broken end showed that the structure was no more crystalline than a similar iron in its natural state, but at the same time appeared to have acquired here and there characters of much interest. . . Instead of all

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Hatfield Memorial Lectures VoL II

the crystals being in close contact all around, some appeared as if slightly separated. It thus appears to me that we may easily understand why repeated shaking and flexure, which bring into play forces in no way adequate to break the bar at once, may yet be able to separate first one crystal and then another, where the strain is at a maximum, until the structure becomes so far disorganised that fracture occurs. Sorby had wondered if the famous Widmannstatten figure, observed in meteorites, could also be found in steels. As recorded in his diaries.l? he observed what he thought was the Widmannstatten figure on 28 July 1863. C. S. Smith.I-' who was the first to draw attention to this entry, has described this as ' ... the very day on which modern metallography was born'. Sorby's earliest work was done with visual magnifications of up to 200 diameters, and he found that in almost every case 'a power of 50 linear showed, on a smaller scale, as much as one of200'. This led him to conclude that he had seen the ultimate structure. He devoted much attention to what he called 'the pearly constituent'. When viewed in the oblique illumination provided by a parabolic reflector of his own design, this constituent 'had the appearance of finest mother of pearl'. Sorby realised that the pearly constituent was in fact a mixture, and from its optical characteristics he deduced that it had a fine lamellar structure. When, in 1885, he obtained magnifications as high as 650 diameters with the vertical illuminator designed by Beck, he appreciated the advantages of high power, for he was able to see things previously hidden from him, and in particular to resolve the structure of the pearly constituent. This constituent is now known as pearlite, though the pearly appearance is rarely seen by the modern metallographer owing to his use of vertical illumination, and of magnifications sufficiently high to resolve the lamellae. Sorby correctly concluded that the soft lamellae consisted of substantially pure iron, and the hard ones of iron carbide. In 1864 Sorby photographed some of his microstructures, and later used them to illustrate his 1887 paper, but the method was limited to very low po\vers and there is no permanent record of what he saw at higher magnifications. Even to show the structure of a Bessemer steel ingot at a magnification of 27 diameters, Sorby reproduced a drawing rather than a photograph, and there is little doubt that early metallographers needed to be draughtsmen of some skill if they wished to record their observations. An interesting collection of drawings of microstructures at high powers was reproduced by Arnold!" in a paper on the influence of carbon on iron published in 1896. He wrote: 'The structures were all drawn from the microscope, when necessary a micrometer being used on correspondingly graduated circles 28 inches in diameter. The drawings were then reproduced by photography to the diameter of the microscopic field. The labour involved in carrying out this process was great, but the results depict the structures with an accuracy unattainable by dierct photography'. The original photographic plates on to which the drawings were reduced for reproduction still exist in Arnold's old department, and an example is shown in Fig. 1.

Metallography - A Hundred Years after Sorby

167

A

100,

I ON Fig. 1

IA a44 ON C' 0-38 RI ES ·18

Annealed 0.38% steel (x 100) (Drawing by J. o. Arnold).

Though there is no permanent record of the microstructures observed by Sorby, we are fortunate in that he bequeathed his original microspecimens (Fig. 2), to the Department of Metallurgy of the University of Sheffield. Thanks to certain characteristics of his technique it is still possible to see some of his specimens just as he prepared them. Following the technique he had developed for rocks, he ground a flat surface on a thin slice of the metal to be examined, attached it to a glass slide and carefully polished and etched the upper surface before covering it with a thin glass slip cemented on with Canada balsam. This prevented rapid deterioration, and small areas of some of the specimens are today still free from corrosion and, apart. from instrumental differences, appear just as Sorby saw them. A sample of decarburised white iron preserved in this way and photographed through the cover glass, is shown in Fig. 3; it is singularly free from scratches and includes a large area of the 'pearly constituent'. In spite of the striking developments in metallography that have taken place since last century, and although there are now powerful new ways of studying the structure of metals, much is still done by essentially the techniques that Sorby developed. Probably the aspect of modern techniques that would surprise him most would be the speed with which metal surfaces can be prepared for microscopic examination; minutes are sufficient

168

Hatfield Memorial Lectures VoL II

Fig. 2

Sorby's original microsections (Preserved in the Department of Metallurgy, University of Sheffield).

for preparations that would have taken him days. He would also be impressed by the quality of the optical equipment now available, and by the ease with which microstructures can be photographed at high powers. Looking through a collection of modern photomicrographs he would notice structures well known to him, but sometimes in different alloy systems or on different scales. We may be sure that he would look for the Widrnannstatten structures seen in the meteorites that first stimulated his interest in metals, and which he recorded photographically (Fig. 4). It is now known that the Widmannstatten structure that so interested Sorby is due to precipitation upon clearly defined crystallographic planes, and that it occurs in many alloys, including some of the Co-Ni-Nb system.lf The process of precipitation at 850°C in an alloy containing 62 at.-%Co, 32 at.-%Ni and 6 at.-%Nb is illustrated in Figs. 5, 6 and 7. Mter a short period (Fig. 5), the precipitate is Widmannstatten in character, though much less bold than in the meteorites, particularly when the magnification is borne in mind. On further ageing at 850°C the precipitate largely transforms to a discontinuous, pearlitic form (Fig. 6), and then redissolves to give once more a Widmannstatten structure (Fig. 7), very similar to that of the Tazewell meteorite. The reason why a Widmannstatten precipitate should be favoured both early and late in the process, yet give place to a discontinuous type at intermediate stages, is by no

Metallography - A Hundred Years after Sorby

169

Fig. 3 Decarburised white iron prepared by Sorby in 1863-1865 (x 500). The structure was photographed through the cover class in 1951 without pre-preparation; it includes pearlite, ferrite, globular carbide, grain boundary cementite and non-metallic inclusions.

Fig. 4

Widmannstatten

structure

in the Tazewell

meteorite

Sorby? and C. Hoole in 1864).

(photographed

by H.

c.

170

Hatfield Memorial Lectures Vol. II

Fig. 5

62:32:6 Co-Ni-Nb

alloy. Aged 30 min at 850°C (x 1000).

means self-evident, and the further study of the phenomenon illustrates two points of considerable importance. The first is of general application, namely, that in any metallurgical investigation full advantage should be taken of every technique that may throw light upon the problem; the second, that as metallography has developed it has become possible to place some observations on a quantitative basis and so greatly extend its scope as a tool of scientific research. Before metallography was born, scientists, in their attempts to understand the varying properties of alloys, subjected them to chemical attack and then analysed any insoluble residues. A modification of this method proved valuable in studying the Cc--Ni+Nb system. The precipitate particles were extracted electrolytically at various stages of ageing, and examined by X-ray diffraction and chemically. For about fifty years now the metallurgist has been able to use X-ray diffraction methods to determine crystal structure, and to find out how the atoms are arranged within the crystals seen in the microscope. Such methods revealed that the precipitates extracted from the Co+Ni+Nb alloy were all Laves phases of the MgZn2 type, irrespective of the mode of precipitation. Chemical analysis, however, showed that the composition of the precipitate changed considerably with ageing time. Niobium occupied almost 60% of the atomic sites early in the precipitation process, but less than 30% after 250 hours at 850°C. This information on the constancy of crystal structure and the variation of chemical composition of the precipitates could not have been obtained by metallography, but metallography could provide far more than a description and classification of the

Metallography - A Hundred Years after Sorby

Fig. 6

62:32:6 Co-Ni-Nb

171

alloy. Aged 4 h at 850°C (x 1000).

morphology of the precipitates. Using the methods of quantitative metallography it was possible to determine the relative proportion of Widmannstatren and of discontinuous precipitate at any time, and thus to expose changes that occurred in the kinetics of precipitation. Quantitative metallography is really an exercise in geometrical probability; systematic measurement must be made on a series of plane sections in such a way that it is possible to deduce what is happening in three dimensions. In the present example it is necessary to know the volume fractions of the alloy that are filled with Widmannstatten and discontinuous precipitate respectively. This information can be obtained most simply by point counting. If a sufficiently large area of the microstructure is covered with a network of points, and if the fraction of these points that coincide with regions of discontinuous precipitation is observed, the resulting point fraction is equal to x, the volume fraction of the matrix that has transformed by discontinuous precipitation. If the precipitation process follows. a rate law of the type x = 1- exp (- btn)

172

Hatfield Memorial Lectures VoL II

Fig. 7

62:32:6 Co-Ni-Nb

alloy. Aged 250 h at 850°C (x 1000).

(where t is the time in hours, and band n are constants for a given reaction), a straight line of slope n should result ifln 1/1 - x is plotted against t on logarithmic scales. The data for the 62:32:6 Co+Ni+Nb alloy plotted in this way are shown in Fig. 8. As a first approximation in experimental points can be represented by three straight lines of slopes n 2.6, 1.25 and - 0.8 respectively. These results, together with those for alloys containing 18-40%Ni and 3.5-S.S%Nb, fit in with the view that the discontinuous precipitation occurs randomly over the grain surfaces. At the first change in slope all possible nucleation sites have been activated and the intermediate line corresponds to growth without simultaneous nucleation. The final stage, represented by the line of negative slope, consists of dissolution of the discontinuous precipitate. There are other ways in which qualitative metallography may be used to obtain more detailed information about the way in which a metallurgical change takes place. Thus, many 'transformations occur by nucleation and growth, and it is desirable to know more about each of these processes. The nature of the problem is illustrated by Figs. 9 and 10 which show two stages of the graphitisation that occurs at 650°C in a normalised carbon

Metallography - A Hundred Years after Sorby

173

10

O,O, '1'0

J~

I

I

~I

10

lOa

_

1000

TIM E , h

Fig. 8

Log/log plot of In l/l-x versus ageing time at 850°C for discontinuous precipitation in a 62:32:6 Co-Ni-Nb alloy.

steel containing aluminium, a phenomenon of some practical importance. 16 The graphite nodules so formed are essentially spherical, and in the micrographs they appear approximately as circles, but it is not possible to tell on inspection whether the diameter of a given circular patch is that of the corresponding graphite nodule or not, since the nodule may not have been cut at its maximum diameter. After seven days at 650°C, the structure is mainly pearlitic and there are relatively few graphite nodules (Fig. 9). After 14 days most of the pearlite has transformed, there are far more graphite nodules, and most of them appear to be larger (Fig. 10). In other words both nucleation and growth must have occured in the intervening period, but it is not easy to separate their effects since the small size of a given graphite patch may be due to recent nucleation or to the fact that the particular graphite sphere was sectioned well away from the equitorial plane. Thanks to a mathematical analysis due to Scheil,"? if sufficient particles are measured it is possible to determine how the separate processes of nucleation and growth have proceeded. Sometimes it is advantageous to combine point counting with Scheil analysis, as in a recent study!" of the effect of hydrogen upon transformation to pearlite. By point counting it was found that hydrogen present to the extent of about 1 cm3/100 g in a 0.4%C:2%Cr steel caused a retardation of about 30% in the isothermal transformation to pearlite at 685°e. Scheil analysis showed that hydrogen did not affect the rate of nucleation or the rate of growth of the pearlite, and that where it retarded the transformation it did so by increasing the length of the incubation period before any nuclei were formed. This effect is illustrated by the results plotted in Fig. 11.

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Hatfield Memorial Lectures VoL II

Fig. 9

O.76%C:O.22%Al steel normalised from 1000°C and annealed for seven days at 650°C. 5% transformed.

Fig. 10

O.76%C:O.22%A1 steel normalised from 1000°C and annealed for 14 days at 650°C. 82% transformed.

Metallography - A Hundred Years after Sorby

Nifroqen

Hydroqen

,-, 6-/:.

II0s

.-.

110

0-0

80 s

175

SO s

s

z

o

...J

o

U

u ~...J a= -c UJ

e, u.

o

SI Z E, S Fig. 11 Scheil plot of the total number of pearlite colonies, N, larger than a given size against the size group S for a O.4%C:2.0Cr steel isothermally transformed at 685°C.

Quantitative metallography is of increasing importance search and is involved in the determination of:

in modern

metallurgical

re-

• grain size; required, for example, in yield and fracture stress studies • spatial distribution of discrete particles; necessary to the understanding of modem dispersion hardened alloys • angles between grains and phases; as used so effectively by C. S. Smith 19 in his classical paper' Grains, phases and interfaces' on the factors determining microstructures. Reliable quantitative results demand large numbers of observations and even with modem aids the work may be tedious and slow. Fortunately qualitative observations are adequate for many purposes. This may be illustrated in terms of the graphitisation of ferritic steels referred to earlier. Quantitative metallorgaphy is invaluable in contributing to a detailed knowledge of the graphitisation process, but qualitative observations may be sufficent to decide the suitability of a given steel for power stations use. If graphite is

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Hatfield Memorial Lectures VoL II

observed after a relatively short time at service temperatures the steel must be rejected, and the amount of graphite and its precise rate of formation are then of secondary importance. The value of qualitative observations and the fact that it is not always necessary to etch metals for microscopic examinations are both illustrated in Figs. 12 and 13. When studying the high temperature deformation or creep of metals it is a common practice to accompany mechanical tests by microscopical examination of polished surfaces, for the changes in topography revealed in this way throw much light on the nature of the deformation processes involved. Thus, in a specimen of niobium stressed at 6000 lb/rn-' at 950°C to give a total creep strain of 20%, the predominant feature was coarse slip within the grains (Fig. 12). A similar specimen (Fig. 13), stressed at 1500 lb /in? at 1100°C to give a total creep strain of 10%, showed very little coarse slip but grain boundary migration was prominent, presumably because the necessary diffusion could occur fairly easily at 11000 but not at 950°C. At an early stage in a research of this kind it would certainly be more important to know of this qualitative difference in the principal mode of plastic deformation than to attempt a quantitative assessment of either phenomenon. Successful metallography depends upon the ability to identify individual constituents in the microstructure, and this is not possible from appearance alone. As we have seen, almost indistinguishable Widmannstattcn structures may be obtained from a meteorite and from a cobalt-nickel-niobium alloy. The metallographer interprets his structures in

e

Fig. 12

Niobium stressed at 6000 Ib/in2 at 950°C. Total creep strain 20% (X 270).

Metallography - A Hundred Years after Sorby

Fig. 13

177

Niobium stressed at 1500 lh/in- at 1100°C. Total creep strain 10% (x 150).

the light of his general knowledge of alloys and of his detailed knowledge of the specimen. If he encounters a new constituent he will, if necessary, carry out subsidiary experiments to enable him to identify it. Once identification has been accomplished it is extended, with care, to any constituents which show similar metallographic characteristics in essentially similar alloys. Until recently, the metallographer has relied mainly upon chemical and X-ray tests, and upon response to polarised light, to help him in his identification, but the position has been radically changed with the development of the electron probe scanning microanalyser.s" Essentially, this instrument allows a fine electron probe to be scanned over a selected area of a microspecimen (Fig. 14). Some of the electrons reflected from the specimen surfaces are collected and with the aid of modern electronic devices are caused to give an electron micrograph of the area examined. Simultaneously, there is an emission of X-rays of various wavelengths characteristic of the elements present in the area irradiated. These are passed through an X-ray spectrometer adjusted so that only the X-rays corresponding to a selected element will be received by the counter tube. The intensity of the X-ray beam, as measured by the counter, is used to control the electron beam intensity in a cathode ray tube which is synchronised with the main electron beam in the probe analyser. Consequently the second display tube shows the distribution of the selected element in the area under examination, and by varying the spectrometer setting the distribution of several elements may be studied in succession. At present the method is not applicable to the lighter elements, of atomic weight less than magnesium, though there are indications that these limitations will soon be removed.

178

Hatfield Memorial Lectures VoL II SCANNtNG ELECTRON BEAM

X-RAY

ELECTRON DETECTOR

SPECT ROMETER & DETECTOR

Fig. 14

Line diagram of electron probe scanning micro analyser.

The electron image depends for contrast upon surface topography and differences in atomic number, and when allowance is made for this the electron image agrees well with the light micrograph of the same area (c£ Fig. 15a and b). This is particularly valuable in microanalysis since it means that, for the area being scanned, it is possible to see, side by side in the two display tubes, both the microstructure and the distribution of the selected element as revealed by the emission of characteristic X-rays.

Fig. 15 Microanalyser pictures of surface crack (x 800): (a) (top left) light micrograph; (b) (top right) electron micrograph seen on display tube of microanalyser; (c) (bottom left) distribution of nickel (Ni, Ka, radiation) as seen on X-ray display tube; (d) (bottom right) distribution of copper (Cu, Ka, radiation) as seen on X-ray display tube.

Metallography - A Hundred Years after Sorby

179

Sometimes the microanalyser can give an almost complete answer to the problem presented to it, as when it showed that the occasional non-metallic inclusions that nucleated graphite in a high purity iron-carbon alloy consisted of titanium, chromium and sulphur. Sometimes the problem is much more complex, and much more information is needed than the microanalyser itself can provide, but always it is extremely helpful to know the distribution of elements in a microstructure with unusual features. Among the complex problems to whose solution the microanalyser has made a useful contribution is included crazy cracking on the surface of annealed sand castings in a nickelchrome-molybdenum steel. Crazy cracking of this kind is often attributed to stress cracking along the strings of non-metallic particles which emerge at the surface of the casting. Very careful industrial metallographic work had shown that these particles are frequently associated with a wedgelike zone, apparently metallic in character, which penetrated along the austenitic grain boundaries. The further information that the microanalyser was able to provide about this problem is illustrated in Fig. 15. In the electron image of a crack (Fig. 1Sb) the unknown metallic constituent is white and the non-metallic particles appear as the adjacent mottled areas. In Fig. 1Sc, light areas represent high concentrations of nickel, since the microanalyser had been set to respond to nickel radiation. The almost exact correspondence between Fig. 15c, and the electron image, Fig. lSb, shows that there is a concentration of nickel in both the metallic and non-metallic constitutents associated with the crack with, if anything, a higher concentration in the non-metallic than in the metallic constituent. When adjusted for copper, Fig. lSd, the microanalyser detects concentrations of this element in both the metallic and non-metallic areas associated with the crack and also a much higher concentration near the casting surface; the latter probably results from preferential oxidation of the iron. Similar photographs obtained with chromium and iron radiations respectively, indicate that chromium tends to concentrate in the non-metallic phase and that areas rich in chromium and nickel appearto be depleted in iron. Other cracks were found to have the same general characteristics, though they differed in detail; for example, sometimes there was no evidence of copper segregation to the metallic phase. The microanalyser study established beyond reasonable doubt that an important degree of segregation has occurred. Solidification-segregation, and local enrichment due to preferential oxidation, may both have contributed, but much further work will be needed to establish the nature of the mechanism responsible for this disturbing phenomenon. It was logical to deal with the microanalyser at this stage for the magnifications involved are those of the light microscope, and, even in producing the electron image, the electrons are used in essentially the same way as light in a metallurgical microscope. The microanalyser does not achieve the high. resolving powers associated with modem electron microscopy, and which have been the objective of metallographers ever since Sorby, for the first time, resolved the lamellae in his pearly constituent at a magnification of 650 diameters. The shorter the wavelength the greater the resolving power possible, and so the realisation of the wave nature of electrons and of the very short wavelengths associated with fast

180

Hatfield Memorial Lectures VoL II

electrons, brought the possibility of greatly increased resolving powers. Electron beams can be focused by axial magnetic fields such as are produced by electromagnetic coils. By using a suitable combination of these it is possible to produce an electron analogue of the type of light microscope that is used to examine transparent specimens in transmitted light. This is the form now taken by the standard high resolution electron microscope, but a tremendous amount of work has been put into the design of magnetic lenses and of their associated circuits to achieve the generally accepted resolving power of loA. This compares with 10 000 A for the light microscope, a thousandfold improvement. This improvement was not obtained all at once, yet in metallurgy, factors other than the microscope limited for some years the resolving power that could be achieved in practice. Early work on metals depended upon the preparation of a thin film replica of the metal surface to be studied, and it was this replica and not the metal itself that was examined in transmission in the electron microscope. It was largely owing to limitations of the earlier replica techniques that progress in the metallurgical field was so slow. Two events completely changed the picture. The first was the development of the carbon film replica technique; the second, the realisation that in spite of earlier views to the contrary, it was possible to use thin films of metal as specimens in the electron microscope. The carbon replica technique is simple, yet provides such a faithful reproduction of a metal surface that only the most modem microscopes can take full advantage of it. Improved contrast results from 'shadowing' the replica with C-Pt or with Au-Pd alloy, and this technique was used to obtain Fig. 16.

Fig. 16

0.67%C; 3%Ni steel transformed at 574°C (x 29000). Carbon replica: shadowed Au/Pd alloy.

Metallography - A Hundred Years after Sorby

181

The 'fingers' of pearlite shown in Fig. 16 have developed by sideways growth from a nodule outside the field of view. An unusual feature is that some of the cementite lamellae are joined, or almost joined, to neighbouring ones. A possible explanation is that in the space between the fingers, into which the cementite is growing, the austenite has been denuded of carbon, and this causes the cementite lamellae to grow into the relatively high carbon zones resulting from the formation of the adjacent ferrite plates. Another abnormal pearlitic structure is shown in Fig. 17. In this steel this type of structure has been observed only in specimens transformed at 535°C, a low temperature which presumably provides a greater 'driving force' for the pearlite reaction. The fact that there are so many circular or ellipsoidal particles of cementite suggests that it is in the fonn of rods rather than plates. The spacing of the cementite is also three or four times that normally observed for a lamellar structure, and this too would indicate a different type of growth.

Fig. 17

0.67%C; 3%Ni steel transformed at 535°C (x 29000). Carbon replica: shadowed Au/Pd alloy.

Lamellar pearlite is normally the easiest growth form, relying as it does upon extension of existing lamellae without fresh nucleation. It may be that the greater 'driving force' available at the low temperature of transformation at which this structure is observed, facilitates the nucleation of new and separate cementite particles so that it is no longer necessary for the cementite to follow the easiest growth form. Although orientation relationships are preserved, individual rods show changes in direction. These are probably due to the growing rod encountering an austenite region already denuded of carbon, so that a change of direction is necessary if growth is to continue. These pearlitic structures were obtained in a study of the variation of interlamellar spacing of pearlite with the temperature of formation.P! The electron microscope is

182

Hatfield Memorial Lectures VoL II

specially suitable for this exercise in quantitative metallography because of the uniformity of the transformed structures and the fine scale of some of the pearlite. An interesting result of the work has been to confirm the idea that there exists a critical temperature above which the nickel concentrates in the ferrite lamellae of the pearlite, while at lower temperatures no such partitioning occurs. Because of its inherent strength, the carbon replica can be removed from relatively rough surfaces and this, coupled with the great depth of focus of the electron microscope, enables it to be used in the study of fractures. In his 1887 paper, Sorby said, 'compared with what can be learnt from good sections, the study of mere fractures teaches very little respecting the ultimate structure ... '. This was true of the technique available to him, but with modern methods the examination of fracture surfaces can be very rewarding. By suitable control of the etching to remove the carbon film from the specimen surface, the replica will carry with it insoluble particles and precipitates previously held in the surface layers of the specimen. This is known as the extraction replica technique; it is applicable both to polished and etched specimens, and to fracture surfaces. It not only reveals structural features reproduced in the carbon replica itself, but also enables the small particles extracted with it to be examined directly and simultaneously in the electron microscope. Metals sometimes fracture prematurely because of the formation of a precipitate on the grain faces. The extraction replica is invaluable in examining such fractures, and the next example shows how it helped to provide an understanding of an industrial problem of some importance.F Under certain circumstances steel castings containing aluminium may fail prematurely with an intergranular fracture. Strong circumstantial evidence had led to the widely held view that the intergranular fracture was due to the precipitation of AlN, but it was not possible to isolate and identify any AlN particles, nor did the hypothesis explain why certain types of steel are specially prone to intergranular fracture while others are more or less immune in spite of high contents of both aluminium and nitrogen. The micrograph (Fig. 18) was obtained from an extraction replica of an area of intergranular fracture in an industrial casting. The rounding areas are features of the fracture surface reproduced in the carbon replica; the darker areas correspond to an extensive thin film, apparently brittle, lifted from the fractured surface and examined directly in the microscope. This film has fragmented to some extent, either when the casting was fractured, or during removal of the replica from the fracture surface. An important feature of the modern electron microscope is that by suitable use of the electromagnetic lenses a selected area of the specimen may be examined by electron diffraction to reveal its crystal structure. When examined in this way the precipitate film of Fig. 18 gave the electron diffraction pattern of Fig. 19. This shows that most of the precipitate was in the fonn of a single crystal, though some of the diffraction spots do not conform to the general hexagonal pattern, so that part at least was in a different orientation from the rest. The diffraction spots are those to be expected from AlN, thus confirming the hypothesis. In other replicas the AlN has shown a dendritic structure, and the resulting diffractions have been arcs rather than spots. Consideration of the fracture topography revealed by the

Metallography - A Hundred Years after Sorby

183

Fig. 18 Extraction replica from fracture surface of steel casing showing intergranular fracture.

Fig. 19

Electron diffraction pattern from precipitate film of Fig. 21.

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Hatfield Memorial Lectures VoL II

replicas has shown that sulphur can influence fracture characteristics, even in the presence of aluminium nitride, and this has led to a satisfactory explantion of the anomalies with which the AlN hypothesis was previously confronted. In high sulphur converter steels, the sulphide particles are so numerous, and provide so many centres from which so-called 'ductile' fracture can be initiated, that ductile fracture absorbs less energy than would intergranular fracture caused by AlN. In clean, high quality steels, on the other hand, there are relatively few sulphide and other particles from which ductile fracture can start, and even small amounts of AlN have large effects in promoting fracture. One of the most recent and far reaching developments in technique in recent years is that of transmission electron microscopy, 23 the direct examination of thin films of metals in the electron microscope. With its aid lattice defects such as dislocations have been seen for the first time and some misconceptions have been corrected. A new and powerful tool has become available for the study of a wide range of metallurgical phenomena. Electropolishing has become of increasing importance in metallography in recent years, and it is largely due to improved methods that undistorted metal films about 2000 A thick and suitable for direct examination can be prepared. Transmission thin film electron microscopy has tremendous potentialities; it provides an alternative approach to problems previously studied with the aid of replicas, but, in addition, because no replica is required, it enables some structural problems to be studied for the first time. The following examples have been chosen to illustrate some of the ways in which the range of electron microscopy has been extended by the introduction of this new technique. Thus, a thin film study of the precipitation of NbC in an 18:10:1 Cr-NiNb austenitic stainless steel-" included Fig. 20.

Fig. 20

18:10:1 Cr-Ni-Nb steel (x 120000). Solution treated for 24 h at 1300°C: not deformed; tempered 3 h at 700°C. Thin film specimen.

Metallography - A Hundred Years after Sorby

185

This micrograph (Fig. 20) includes many features that throw light on the way precipitates interact with lattice defects. There is clear evidence of precipitation on dislocations, as well as signs that, as the precipitate particles grow and lose coherency with the matrix, the forces binding them to the dislocation on which they were formed disappear, and are replaced by forces or repulsion that push the disclocation away. In the centre of the field is a region of stacking fault between two partial dislocations. Stacking faults are so called because, in fcc structures such as austenite, whenever a dislocation dissociates into two partial dislocations, the atoms in the region between them are rearranged so that the close packed planes are stacked in the manner required for a cph structure. Close examination of Fig. 20 shows that precipitation is occurring on the stacking fault and causing fringes at an angle to those due to the stacking fault itself This is shown more clearly in Fig. ·21 which corresponds to a later stage in the precipitation process when the stacking faults have extended throughout the matrix.

Fig. 21

18:10:1 Cr-Ni-Nb steel (x 200000). Solution treated for 24 h at 1300°C: not deformed; tempered 5 h at 700°C. Thin film specimen.

The stacking fault fringes are now fully resolved and they extend beyond the boundaries of the micrograph. As the precipitate particles lose coherency, the stacking faults begin to be eliminated as can be seen from the two regions which contain no stacking fault fringes, and in which the precipitate particles are clearly visible. Contrast in micrographs obtained from replicas is due mainly to the varying thickness

of replica materials traversed by the electrons. In thin film electron microscopy the

186

Hatfield Memorial Lectures VoL II

principal source of contrast is Bragg reflection of the electrons. It is therefore necessary to be sure that every feature of a thin film micrograph can be interpreted both metallurgically and in terms of diffraction effects. There are those who take the view that the central features of Figs. 20 and 21 are trains of closely spaced dislocations, rather than stacking faults, but such an hypothesis cannot account for all the features observed. Diffraction contrast can be turned to good account in identifying constituents, as illustrated by Figs. 22 and 23, which are photographs of the same area of an averaged specimen of the 18:10:1, Cr-Ni-Nb steel.

Fig. 22

18:10:1 Cr-Ni-Nb steel (x 120000). Solution treated for 24 h at 1300°C: not deformed; tempered 24 h at 700°C and 4 h at 850°C. Thin film specimen.

Fig. 23

18:10:1 Cr-Ni-Nb

steel (x 120000). Solution treated for 24 h at 1300°C: not

deformed; tempered 24 h at 700°C and 4 h at 850°C. Dark field illuminated.

Metallography - A Hundred Years after Sorby

187

In Fig. 22, the bright field image obtained in the normal manner shows mainly well developed precipitate particles on (111) planes in various orientations and on individual dislocations. Dark field illumination is achieved by moving the objective aperture so that only a chosen Bragg reflection can contribute to the image. For Fig. 23 one of the Bragg reflections of NbC was chosen, with the result that all Bragg reflections arising from the austenite matrix were excluded. Under these conditions there are certain lens aberrations and this causes a blurring of the image, but it is nevertheless clear that there is almost perfect correspondence between the precipitate particles as revealed by the two methods of illumination, and that therefore they are all of the same kind whether they occur at stacking faults or on dislocations. This method can be used to distinguish particles about 30 to 50 A across, separated by about 100 A. For example, with a thin film containing carbides of both Ti and U, if a TiC reflection is used the U C particles do not show up, and vice versa. Sub-grain formation in ferrites, or alpha veining, has been a subject of interest since the early days of metallography and much work over the years has enabled its nature to be elucidated. A single thin film micrograph.s> (Fig. 24) illustrates many of the accepted ideas.

Fig.24

0.2%C: 4%Mo steel (x 160000). Quenched and tempered 1000 h at 550°C. Thin film specimen.

This high magnification micrograph shows sub-grain formation in a precipitate free area of a quenched and heavily tempered fum of a 4%Mo steel. The sub-grain boundaries are clearly formed from dislocation networks, and there is contrast between adjacent subgrains because of the difference in orientation between them.

188

Hatfield Memorial Lectures Vol. II

Transmission electron microscopy may also have a contribution to make in the study of magnetism. If a thin film, ferromagnetic specimen is defocused in a region of low magnetic field, the domain walls appear as white or black lines.?" as in Fig. 25, and dislocation tangles appear as diffuse black areas. The domain walls show little interaction with the dislocations.

Fig. 25

Thin film of iron (x 50000).

It would be possible to list many more applications of the thin ftlm technique, but in spite of the success already achieved, probably the most exciting ones still remain in the future. With the development of hot and cold stages for electron microscopes, it becomes possible to do experiments inside them. These may include the systematic study of metallurgical processes at various temperatures, such as precipitation, graphitisation, recrystallisation, oxidation and the growth of films, to mention but a few. There is a general desire to extend the technique in this way as soon as the additional facilities are available and it is possible to reserve a microscope for a relatively long period for a single experiment. Electron microscopy has so greatly extended the range of metallography, and has so many exciting potentialities, that it is easy to give the impression that the light microscope is now outdated, so that it is only a matter of time before it will be completely superseded by the electron microscope. To do so would be quite wrong, for many important metallurgical phenomena are adequately resolved by the modern light microscope; it is for this reason that light micrographs are chosen for the final examples, Sorby wrote: 'The changes of structure produced by hardening deserve far more study, but will, I fear tax to the uttermost the capabilities of the microscope, since the constituent grains

of the hardened steel are so extremely minute'.

Metallography - A Hundred Years after Sorby

189

This view is supported by a drawing of the martensite in a hardened steel made a few years later by Arnold.!+ and reproduced in Fig. 26.

H

600 IRON

CARBON C~ IMPURITIES Fig. 26

DIA 98···92 0·89 0··19

0.89%C steel in hardened condition (Drawing by J. o. Arnold-+) (x 600).

The scale of the martensitic structure can be increased by using coarse grained highly alloyed steels, and by hardening at sub-zero temperatures. Figure 27 is a light micrograph obtained in this way. The specimen showed a sudden 'burst' of transformation. The morphology is that commonly associated with martensite transformation and, particularly in the central area, there are certain features that are observed only ~hen transformation occurs in 'bursts'. When transformed in liquid nitrogen (- 196°C), and examined at the highest magnification possible with the light microscope, this steel shows features that have only recently been discovered by electron microscopy. Figure 28 shows fine martensite needles branching from a main martensite plate with internal features within this plate clearly resolved. An internal structure, either slip or twinning, was predicted theoretically.i" and later observed and shown to be twinning by electron microscopy.P? The fact that such fine detail can be resolved in the light microscope is an indication of the way in which conventional metallography has responded to

190

Hatfield Memorial Lectures VoL II

Fig.27

O.51%C:24%Ni steel (x 145). Transformed to martensite by cooling to -80°C.

the stimulus of the electron microscope, and a sure sign that the light microscope still has an important part to play in the study of metals. So ends this survey of metallography a hundred years after Sorby. In the time available, it has not been possible to refer to all the techniques that properly come under this heading. I have preferred to take my examples from techniques actively pursued in a single department (my own) and I have excluded polarised light microscopy, phase contrast, interference techniques, autoradiography, microradiography, and field emission microscopy. Even within the field covered, there has been the need to be highly selective and since selection is essentially personal, others might have chosen differently. However, I hope that sufficient has been said to demonstrate the great and still growing importance of metallography, and to justify our paying tribute to Henry Clifton Sorby as the father of the subject.

ACKNOWLEDGEMrnNTS l'1y grateful thanks are due, and gladly given, to all my colleagues for help In the

preparation of this lecture, particularly by providing the illustrations. I am also grateful to

Metallography - A Hundred Years after Sorby

Fig.28

191

0.51%C: 24%Ni steel (x 2200). Transformed to martensite by cooling to -196°C. Light micrograph.

Mr H. R. Singleton, Director of the City Museum; Sheffield, for allowing me to reproduce the portrait of Dr Sorby.

REFERENCES 1. C. H. DESCH: 'The Services of Henry Clifton Sorby to Metallurgy', second Sorby Lecture, Sheffield, 1921. The Sorby Lectureship was instituted '. . . to commemorate the work of Dr Henry Clifton Sorby, FRS, who rendered such signal service to science in general, and to the special branch of microscopy in particular'. Professor W. G. Feamsides (first Sorby Professor of Geology) delivered the first Sorby Lecture on 28 February 1914, before the Sheffield Society of Engineers and Metallurgists. Later lectures were organised by a joint committee of that Society with the Sheffield Association of Metallurgists and Metallurgical Chemists (later the Sheffield Metallurgical Association), the Sheffield Local Section of the Institute of Metals, the Sheffield Branch of The Institution of British Foundrymen, the Sheffield Section of the

192

Hatfield Memorial Lectures VoL II

Junior Institution of Engineers, and the Sorby Scientific Society. Lectures were given in 1921, 1923, 1926, 1928 and 1930, and the lecturers included Walter Rosenhain, H. C. H. Carpenter and F. C. Thompson. The lectures were published and sold for a nominal sum, but were discontinued in 1933, for lack of financial support. 2. C. S. SMITH: A History of Metallography, University of Chicago Press, Chicago, 1960, 168185. 3. ROBERT HOOKE: Micrographia, London, 1665. 4. N. T. BELAIEW:Rev. Met., 1914, 11,221-227. Belaiew called attention to the fact that a Russian, P. Anasoff, tried in 1841 to imitate the watered pattern of Damascus steel and used the microscope to examine polished and etched steel surfaces. Anaso£rs paper was in Russian and remained unknown outside Russia until the early years of this century. 5. L. S. BEALE:How to Work With a Microscope, 4th edn, London, 1868, 181-183. 6. A. MARTENS: ZVDI, 1878,22, 11,205,481; 1880,24,398. 7. H. WEDDING:JISI, 1885 (i), 187-199. 8. H. C. SOREY: 'The Structure of Iron and Steel', (Abstract),JISI, 1882 (ii), 702-703. 9. H. C. SOREY: 'On the Microscopical Structure of Iron and Steel' ,JISI, 1887 (i), 255-288. A preprint of the same title was issued for the 1885 Annual General Meeting of the Iron and Steel Institute, and was essentially a summary of the 1887 paper. It is item No. 79 in Volume II of Sorby' s collected works in the library of the University of Sheffield, and is reprinted as an appendix to Ref 2. 10. H. C. SOREY: 'On the Application of Very High Powers to the Study of the Microscopial Structure of Steel' ,JISI, 1886 (i), 140-147. 11. ]. O. ARNOLD:]. Sheffield Technical School Metallurgical Society; account of meeting held 31 October 1891. 12. H. C. SOREY: 'Diary', 1859-1908 (incomplete), University of Sheffield Library; available on microfilm from Micromethods Ltd., Wakefield. 13. C. S. SMITH: A History of Metallography, University of Chicago Press, Chicago, 1960, 172. 14. J. O. ARNOLD: Proc. Inst. Civ. Eng., 1896, 123, 127. 15. B.]. PIEARCEYet al.:JIM, 1962-3,91,257. 16. ]. E. HARRIS et al.: 'Steels for reactor pressure circuits', lSI Spec. Rep. 69, 1961,54-76. 17. E. SCHElL: Z. Metallk., 1935,27,199-209. 18. ]. H. WOODHEAD: Private communication. 19. C. S. SMITH: Trans. AIME, 1948,175, 15. 20. P. DUNCOMB: Brit.]. App. Phys., 1959,10,420-427. 21. R. BOOTH and J. H. WOODHEAD: unpublished work. 22. J. A. WRIGHT and A. G. QUARRELL:JISI, 1962,200,299-307. 23. P. B. H. HIRSCH: Met. Rev., 1959,4,101. A. HOWIE: Met. Rev., 1961, 6, 467. 24 .. R. W. K. HONEYCOMBE et al.: N.P.L. conference on structure and Strength of Alloys, January 1963. 25. J. J. IRANI: PRIVATECOMMUNICATION. 26. D. H. WARRRINGTON: Private communication. 27. R. BROOK and A. R. ENTWISTLE:Private communication. 28. M. S. WECHSLENet al.: Trans. AIME, 1953, 197,1503. 29. P. M. KELLYand]. NUTTING: Proc. ROy. Soc., 1960, 259A 45-48.

SEVENTEENTH

HATFIELD

MEMORIAL

LECTURE

Interrnetallic Chemistry of Iron w.

Hunte-Rothery

At the time the lecture was given Professor Hume- Rothery was Isaac Wolfson Professor of Metallurgy at the University of Oxford. The lecture was presented at the Firth Hall of the University of Sheffield on Wednesday 19 May 1965.

At this time of year we pay tribute to one of Sheffield's most distinguished scientists, Dr William Herbert Hatfield who entered what was then the University College of Sheffield at the turn of the century. He was a Sheffield man from beginning to end, with a great pride in his home town and university. His scientific work is well known, and earned for him the Fellowship of the Royal Society at a time when very few industrial scientists were elected. It is a high distinction to be asked to lecture in honour of such a man, and I can only thank you for your kindness in inviting me today. In choosing a subject, I have borne in mind that most of Hatfield's work was in connexion with alloy steels. The years of Hatfield's work saw amazing advances in the science and understanding of the subject. He was full of enthusiasm for the new methods, and I can remember his active participation in meetings when X-ray metallography was being discussed. Since his death the new work has proceeded with ever increasing speed, and it is perhaps useful to stop for a moment and try to survey one section of the whole. Owing to their complexity, I shallnot deal with the actual alloysof industry but shall try to answer the more simple question 'What sort of alloy will iron form with each element of the periodic table?'. Hatfield would have approved such an enquiry, for he realised that the simple foundation must be laid before the complicated superstructure could be built.

Our problem today is, therefore, to take the periodic table as a whole (Fig. 1) and to examine the extent to which we can generalise or interpret the structures of the alloys which iron forms with different elements. Pure iron exists in the body centred cubic form at high (B-Fe) and low (a-Pe) temperatures, while the ..face centred cubic form (y-Fe) is stable over the range 910°C (A3) to 1389°C (A4)' The curious reversal of phase changes at the A3 and A4 points is regarded as due to magnetic effects, and it is unlikely that any simple theory will explain the effects of different elements ,on the A3 and A4 transformation. In spite of much theory and speculation there is no satisfactory electron theory of the iron crystal. We are justified in regarding the structures (Fig. 2) as those of reasonably hard spheres in contact, and we may regard the surfaces of the spheres as diffuse, with

193

194

Hatfield Memorial Lectures VoL II H~ 2

H I

/ / /

/'

/

"-

""

"-

/

'

lill1rli K

Co

19 2[

P.b 37

cs 15\

SC Ti

V

21 22 2l

Cr

Mn Fe Co Ni

24 25

Zn Go

26 27 28 29 30 31

-. ir I

~

As Se Br

Kr

34 35

~~~\\ \~\\~\\ ~\ So La ~ 'S7 ~I

Hf To W ~

Os

Ir

Pt

All

Hq 11

72 73 74 75 16 TT 18 79 80

Fr P.o Ac :"Th-: Po U Np 87 . 88 89 :_~_: 91 92 93

Fig. 1

Cli

Am 94 95

Pli

em 96

Pb

Be Po At

Rn

81 82 83 84 85 86

Sir. Cf 97 98

Periodic table of elements. From: W. Hume-Rothery and H. Raynor; The Structure oj Metals and Alloys, The Institute of Metals, London, 1962.

electrons in hybrid spd orbitals boiling over from one atom to the next and holding the structure together. The closest distances of approach of the atoms in these structures are a-Fe 2.48 kX and y-Fe 2.57 kX (at 916°C). The atomic diameter is thus about 2.5 kX, and this is a fundamental value, of great importance in understanding iron alloys. By the size factor of a solute we mean the difference between the atomic diameters of solvent and solute expressed as a percentage of the former. The size factor is the first factor required to understand alloys of iron, and we take the atomic diameters to be the closest distances of approach of the atoms in the crystals of the elements. We then need some quantity to express the effects which are found when one metal in an alloy is very electropositive compared with the other. For this purpose the electronegativities of Pauling may be used and are shown in Fig. 3, which shows that iron and silicon have the same value on the electronegativity scale.

Intermetallic Chemistry of Iron

195

Fig.2 The structure offcc(y) iron and bee (a8) iron. From: W. Hume-Rothery and H. Raynor; The Structure of Metals and Alloys, The Institute of Metals, London, 1962. It is well known that in many alloy systems, structure variations follow clear valency principles. Iron shows several valencies, and there is no agreement as to the valency in metallic iron. When dealing with transitional metal alloys I shall, therefore, speak of group number effects, rather than valency effects, and I shall usethe·woup numbers

1

2

3

K Rb Cs

Ca Sr Ba

Sc

Y

4 .Ti Zr

La

HE

5 V Nb Ta

6 Cr Mo W

7 Mn Tc Re·,

8 Fe Ru Os

9 Co Rh Ir

10 Ni

Pd Pt

An equiatomic alloy of iron and chromium has, thus, an average group number (AGN) value of7. In order to understand the alloys of iron, we require to know the effects of different elements upon the melting point of iron, and on the A3 and A4 transformations. As regards the latter, it is well known that solute elements divide themselves into two classes: (i) ferrite stabilisers, which give rise to equilibrium contracted y-fields (Fig. 4a and b)

diagrams with closed y-loops or

196

Hatfield Memorial Lectures VoL II Xp 5~--------------------------------------------------------------~

• F

Fig. 3

I

1

1

1

I

!O

20

30

40

50

I S-f

I

I 71

eo

I

The electronegativities of the elements (Pauling). From: W. Hume-Rothery: Structure of Alloys of Iron: an Elementary Introduction, Pergamon, Oxford, 1966. CLASS

CLASS

I

The

n

M

«

Sa

5b

L...-

4b

Figs. 4 and 5 4a and b: Closed y-Ioop and contracted y-field equilibrium diagrams. Sa and b: expanded and open y-field equilibrium diagrams. From: W. Hume-Rothery and H. Raynor; The Structure of Metals and Alloys, The Institute of Metals, London, 1962.

Intermetallic Chemistry of Iron

197

(ii) austenite stabilisers, which give rise to expanded y-fields (Fig. Sa and b), or to open y-fields in some systems where the second element has an fcc structure Just as the A3 and A4 points may be either raised or lowered, according to the nature of the solute, so the melting points (solidus curves) and the associated liquidus curves may be either lowered or raised. In the alloys with austenite stabilisers we encounter liquidus curves for both 0- and y-Fe, connected by a peritectic horizontal.. It is, therefore, convenient to interpret these alloys in terms of diagrams such as those of Fig. 6a and b in which 1528°C is the hypothetical melting point of "{-iron, and the figures refer to the

lowering and raising of the liquidus respectively.

15360 15280

Fig. 6

Peritectic type of diagram.

198

Hatfield Memorial Lectures Vol. II

We have now to ask to what extent we can understand the effects of different elements on the liquidus and solidus curves, on the A3 and A4 points, on the limits of the solid solutions, and on the existence of intermediate phases. There is; as yet, no theory which enables us to calculate these effects from first principles, but clear general principles have been established empirically. Our first generalisation of the structures of iron alloys is in terms of the size factor, and the size factor principle states that, if the atomic diameters of solvent and solute differ by more than about 14-15%, substantial solid solutions will be restricted. It is important to realise that the principle is a negative one which enables us to say when wide solid solutions are not likely to be formed: it does not enable us to say that they will be formed. Figure 7 shows the atomic diameters of the elements, and the dotted lines, which are drawn at distances of + and -15% of the value for iron, mark out the zone of favourable size factor. This diagram indicates at once the elements which are unlikely to fonn appreciable solid solutions in iron. It shows that the whole series of elements from V-Cr---7Ge---7 As is of favourable size factor but that, on passing back to titanium, the boundary of the favourable zone is being approached. In the second and third transition series the atoms are a size larger, and it is the Group V elements, niobium and tantalum, which are near the borderline. Boron occupies a curious position in which the atomic diameter is too small for substitutional, and too large for interstitial solid solution in iron. Its solid solubility is very

Or-----.----------------::e::-----------------r.6

~

~

•

Yb

e

No 1·4

0

35 -

.A

1

.~ I.

M~ se ~ 3.0~Ci----------------------

,---:-----~--~---~-;

~.

f'

z~

• Mn Zn V Co. • eA.8Cu X As Cr Fe Ni Go Ge

u 2·5

l:

2

<:

X

xt

• A

Si

Se

-~~----~--------------------------------X S

XB C

'·51--..:..t--------------------------------10·6

Fig. 7 The atomic diameters of the elements; the dotted lines mark out the zones of favourable size factor for alloys with iron: the upper broken line represents the ideal value for Laves phases of the type FezM. From: W. Hume- Rothery: The Structure of Alloys of Iron: an Elementary Introduction, Pergamon, Oxford, 1966.

Intermetallic Chemistry of Iron

199

small, and most of the boron atoms are probably accommodated in dislocations, grain boundaries, or other imperfect parts of the crystal, and it is these atoms which produce such remarkable effects on the rates of transformation of steels. Carbon and nitrogen have atoms favourable for interstitial rather than substitutional solid solution. The same would apply to oxygen if it formed metallic or covalent bonds with iron, but actually it is so electronegative that it forms divalent 0- =ions, which are large, and the solubility of oxygen in iron is small, partly for this reason and partly because of the existence of very stable oxides which are formed at the expense of the solid solution, in accordance with the general free energy principle. The size factor also gives some indication of the systems in which Laves phases may be expected. The ideal radius ratio for these is 1.225, but recent work has shown that in transitional metal alloys Laves phases may be formed at radius ratios varying from 1.06 to 1.40. The condition is, thus, not a very stringent one, and we can only say that there is a probability of a Laves phase existing when the radius ratio is between 1.15 and 1.35. We expect a greater variation for the MgZn2 and MgNi2 structures with variable axial ratios than for MgCu2 where the atomic positions are fixed. We may now ask ourselves the question: what will be the effect of a given transitional metal upon the freezing points and melting points of iron alloys? The simplest fonn of equilibrium diagram is that of Fig. 8a where the two metals A and B have the same crystal structure, and the size factors, electrochemical factors, etc. are so favourable that the liquidus and solidus rise smoothly from A to B and a continuous solid solution is formed. The higher melting point of B means that the strength of cohesion in the solid relative to that in the liquid is greater in B than in A. As the atoms of A are replaced by those ofB, these characteristics are carried over into the solutions, whose melting points rise. We may call this the melting point or mp effect, and we may expect to see it in all systems where A and B have the same crystal structure, even though intermediate phases are present. Where A and B are of different crystal structure the position is more complicated, but if A and Bare of the same nature (e.g. ifboth are transitional metals with spd bonding), we

B

A

d.p

B

A

A

solii solution d.~

o

100

a

Fig. 8

B

0 b

solid

solution

100

Three simple types of liquidus and solidus curve.

200

Hatfield Memorial Lectures VoL II

may hope to see the results of the mp effect in a general way, even though no exact correspondence with the actual melting points can be expected. We shall not expect to see the effect so clearly where the two metals differ so widely that it is very improbable for one to assume the crystal structure of the other. We may now consider what may be expected when the atomic diameter of B becomes increasingly different from that of A. The less regular structure of the liquid solution can accommodate stranger atoms more easily than the solid solution, and this tends to favour the liquid phase and to depress the liquidus and solidus curves of both A rich and B rich alloys.At the same time, as atoms of A are replaced by those ofB, the mp effect makes itselffelt, and the liquidus and solidus curves may pass through a minimum as in Fig. 8b, which may change into that of Fig. 8c as the size factor or other factors make complete miscibility impossible. For b-Fe there are no examples of equilibrium diagrams of the type of Fig. 8a. There are, however, examples of minima in the liquidus and solidus. Figure 9 shows the curves for the systems Fe-Cr, Pe--Mo and Fe-W. The size factor effect is responsible for the steeper fall of liquidus and solidus in Fe-Mo as compared to Fe-Cr. The very high melting point of tungsten means that the mp effect is here much more powerful and is responsible for the shift in position of the minimum.

I

Fig.9

sso

Liquidus and solidus curves for the systems Fe-Cr, Fe-Mo and Fe-W.

For "{-Fe,the system Fe-Pt (Fig. 10) is an example of the type of Fig. 8a. In the system Fe-Pd (Fig. 11) the mp effect is very small and so the size factor plays a more important part and we obtain a minimum in the curve. In the Fe-Ni system (Fig. 12) the size factor is very small and the depression of liquidus and solidus is correspondingly smaller than for Fe-Pd; there is a very slight minimum.

Intermetallic Chemistry of Iron IBOO

r--------------------.

Fe Pt

I y Fe] Pt

I

1000 A3

\

~

/'\

\\ \ \

\ \

\\w\

\

SO

100

Pt.ot~O/o

The system Fe-Pt.

Fig. 10 1600

r-----------

1500

u o w

a:.

Fe Pd

::>

I

~JOOO w

0..

~

w I-

....

_-

"'"

I

•....

I

,, \

/ / \

/

I

\

\

500

I I I / I I I /

I

I

\\'f\.

o

50 Pd. or-

Fig. 11

100 %

The system Fe-Pd.

201

202

Hatfield Memorial Lectures VoL II

o

20

400

60

NICKEL ,at-Io

Fig. 12

The system Pe--Ni. From: W. Hume-Rothery:

The Structure of Alloys of Iron: an Elementary Introduction, Pergamon, Oxford, 1966.

We now want to ask whether these effects can be put more quantitatively for different solutes. To understand this, we may consider typical
T

ATOMIC

Fig. 13

Xs

FRACTION

and

Xe

OF SOLUTE

on liquidus and solidus curves.

Intermetallic Chemistry of Iron

203

Figure 14 shows the values of L1G obtained at Oxford for the effects of the transition metals on the b~liquid transformation. In the first transition series manganese behaves abnormally, but there is a clear general tendency to form a minimum in the region of group VII. The points in this figure place the elements in the order of their liquid stabilising powers, in the sense that a large positive value of L1G favours the liquid state. These values include both the mp effect and other effects such as those due to atomic diameters, electronic effects, etc.

Fig. 14

Solute free energy differences ~G for the

B to liquid transformation.

We may therefore write L1G

= f(mp) + L1F

where L1F represents the part of L1G that is not due to the mp effect. Buckley and HumeRothery- have shown how fimp) can be estimated approximately, and in this way it is possible to obtain a rough estimate of L1F, the results being shown in Fig. 15. The point for zirconium may be in error owing to a very steeply falling solidus. For solutes from the first long period, the L1F values are systematically smaller than those for the later periods, and this is what would be expected from the size factors. The point for manganese is now in sequence, and there is a clear group number principle for the first long period. Except for zirconium, the values for corresponding elements in the second and third long periods are reasonably similar, as would be expected from the atomic radii. To summarise: the depression or elevation of the solidus and liquidus curves of ()-Fe are controlled (a) by the melting point effect, and (b) by the size factor and electronic effects. The latter are summarised in Fig. 15, and the total effect in Fig. 14. We can thus begin to understand the form of this part of the equilibrium diagram. We have seen that the A3 and A4 points may be eitherraised or lowered according to the nature of the solute. These curves may be analysed in the same way that the liquidus

204

Hatfield Memorial Lectures VoL II

u?-3 Au


VI

Fig. 15

VII VillA ViliS Ville GROUP No.

8.F values for the

B to

IB

liquid transformation.

and solidus curves were analysed above. As the A3 points are subject to considerable temperature hysteresis, it has been found more convenient to determine the A4 points. This has been done by the Oxford workers, and their results permit the calculation of L1G, the free energy of transference of solute for the iJpy transformation; the convention used is that a positive sign indicates a ferrite stabiliser. Figure 16 shows the results obtained. The result for manganese in the first long period is anomalous, but for the

TI

• • ,

~

I (;

,

bee fee cph

Fovourinq

elements

0
GROUP

Fig. 16

8.G values for

B~y

A 8 C '--VIII-----l No.

18

transformation; a positive value of 8.G indicates a ferrite stabiliser.

Intermetallic Chemistry of Iron

205

remaining transition metals the results show that elements to the left of Group VIII are ferrite stabilisers. There is a clear group number effect except that, as for the Iiquidzzsolid transformations, zirconium has a smaller effect than would be expected from its position in the periodic table. The transition metal austenite stabilisers form a compact group of elements in Groups VIllA, Band C, copper and gold in Group lA, and manganese in Group VIlA. The y-stabilising powers of osmium, iridium and platinum are particularly high. When we go outside the transitional metals, we find that for substitutional solutes most if not all elements are ferrite stabilisers, which interstitial solutes are austenite stabilisers. Taken as a whole, therefore, substitutional austenite stabilisers are in the minority, and are confined to one small part of the periodic table. The effect of interstitial solutes in stabilising austenite can readily be understood, because the holes between the atoms in the fcc structure are larger than those in the bcc. In spite of much speculation, there is as yet not quantitative theory and little qualitative understanding of the effects of the substitutional solutes. The fact that vanadium, niobium, tantalum, chromium, molybdenum and tungsten, which have bee structures of great stability, are ferrite stabilisers might suggest that their atoms possessed electronic configurations which favoured the body. centred coordination so strongly that this characteristic was carried over into their alloys. This may well be true for these elements, but even so we do not understand why the bee ferrite is favoured by aluminium (fcc), beryllium (cph) and silicon (diamond type of structure). At present, therefore, there is a clear empirical relation but no real understanding, and experiment is ahead of theory. Weare now in a position to begin our equilibrium diagrams in the form of Fig. 17. How far are the limits of solubility likely to extend, and what other phases may appear? In answering this question we find again that manganese is anomalous, and must be omitted from our generalisations. Let us consider first the y-loop kind of diagram formed by elements to the left of iron in the periodic table. We find that the y-phase does not extend very far into the equilibrium diagram, and its extent appears to be governed by group number principles. To show this we must first draw our equilibrium diagrams in terms of average group numbers. For the solid solutions in fcc iron, cobalt and nickel the phase boundaries appear as in Fig. 18. This suggests clearly that the fcc phase becomes unstable when the AGN is less than about 7.7. It is very remarkable that the limits of the y-Ioops are at roughly the same AGN values as the limits of the solid solutions of vanadium and chromium in cobalt or nickel. But manganese does not follow this principle. The above relation holds only when solvent and solute are in the same period. When we consider alloys of iron with the elements from the second and third long periods we have solutes whose atoms are one size larger than those of the corresponding elements in the first long period. Since the bee structure is less tightly packed, we shall expect it to accommodate these larger atoms more readily than does the fcc structure, i.e. the bec structure will be favoured and the y-Ioop will be narrowed. This is confirmed by the following values

206

Hatfield Memorial Lectures VoL II

1500

A.•

~~ ~

/",

~

~,-,

~

//

~,

"

u

0 ••

UJ

a::

::>

'< a: UJ

Q...

~ ~

UJ

JOOO

Fig. 17

Left hand side of an equilibrium diagram.

Maximum extent ofy-Ioop (at.-%) Ti Zr Hf

0.6 small ?

V Nb Ta

1.5 1.2 0.95

Cr Mo W

13 1.5 1.0

Continuing with the y-Ioop type of diagram, we may ask how far the aB solid solution will extend into the diagram. In the first transitional series, we have seen that titanium is on the borderline, while vanadium and chromium are within the zone of favourable size factor. Here we find continuous bee solid solutions in the systems Pe--Cr and Fe-V, but in the system Fe- Ti we have entered the region favourable for Laves phases (radius ratio 1.16) and there is a very stable Laves phase, Fe2Ti (Fig. 19).·The solubility of titanium in aB Fe is thus restricted first by the borderline size factor and secondly by the usual free energy effect of an adjacent stable phase. On passing to zirconium in the next period, the solute atom is a size larger and the aB solid solution is severely restricted, but the radius ratio (1.27) is still appropriate to the Laves phase and Fe2Zr is formed (Fig. 20).

Intermetallic Chemistry of Iron

207

1500

1300

Ltq~nd

1100

__ •..•. x-x-x-x 0-0-0-0

-•••••

Fe-V Fe-Cr

Co-v

Co-Cr Ni-V Ni-Cr

900

8'0

7'0

AVERAGE

Fig. 18

9{) GR.OUP

10-0

NUMBER

Solid solution limits of fcc Ni, Co and Fe phases plotted in terms of average group number. From: W. Hume-Rothery: Philos. Mag., 1961,6,769.

The size factor diagram shows that, in the second and third long periods, it is the Group V elements niobium and tantalum which are on the borderline of the favourable zone, and as their radius ratios with respect to iron are about the same as that of titanium we find Laves phases Fe2Nb and Fe2Ta, and there is a general resemblance between and equilibrium diagrams of the systems Fe- Ti, Fe-Nb and Fe- Ta in the iron rich regions (Fig. 21). We have considered how far the bee aD phase extends when iron is alloyed with elements lying to the left in the periodic table. Can we now say what will happen when we alloy with elements to the right of iron in the periodic table, and in particular with the B sub-group elements of Fig. 22? In this figure magnesium is marked with a circle because it does not alloy with iron for reasons we shall consider later. We may consider for a moment the later elements in the second and third long periods. Here we have a sequence of crystal structures: bcc~(j~cph~fcc Groups V and VI bcc Nb, Ta,Mo,

W

Groups VII and VIllA cph Tc, Ru, Re,

as

Groups VIIIB andC fcc Rh, Pd, Ir, Pt

208

Hatfield Memorial Lectures VoL II

Fe Ti

I

~

1500

y

UJ

a: ::>

glooo

1000

c,

L UJ

I-

500

Fe Ti~

Fig.19

I 9(X)

10 20 30

0.

500

Tt

The system Fe-Ti.

40

50

Zr, wt-Ofo 00

it)

80 85 90 95

.E N I

I BOO 1700 1600

I

"

" I'"

I'

0/ !

,P

, •• I

·

I I

,

I

I

0,' I

I

, I

"

I I

I

I---x~x-,,-\ 0

.:~~~j 6000

Fe

Fig.20

10 -20

30

0

\~\/

0

~~

Zr

---------0.-;' \

4'0

SO 6rJ7o-

Zr.ot-%

80

90

,

100

Zr

The system Fe-Zr. From: W. Hansen: Constitution of Binary Alloys, McGra\v-Hill, New York, NY, 1958.

Intermetallic Chemistry of Iron

209

1000

Ti,ot-'.

Fig. 21

The systems Fe- Ti, Fe-Nb and Fe- Ta. From: W. Hansen: Constitution of Binary Alloys, McGraw-Hill, New York, NY, 1958.

rrElJ-

Co -

Ni -

Fig. 22

Cu -

@-

AI -

Si

-

p -

S

Zn -

Go -

Ge

-

As -

Se

B sub-group elements.

Here we find that if we alloy, say, molybdenum with rhodium we may obtain a cph alloy if we adjust the composition to give an AGN of about 7. We have, in fact, a series of alloys in which the above three types of crystal structure, and also the (J' structure, each

appear to be stable over a characteristic range of AGN values, and parts of the equilibrium

210

Hatfield Memorial Lectures VoL II

diagrams are roughly superposed if we draw them in terms of AGN values. Figure 23 shows to a very rough approximation the composition limits of the different phases in terms of AGN. This is a drastic oversimplification but it shows the general tendency. In the first long period, the cph structure drops out and everything is much less clear cut. Figure 24 shows the equilibrium diagram of the system Fe-Co, and this suggests

3000

2500

2000

\

\ \

Mo-Pd

}- -

Ru -Pd

/

1500

1200~--~rO~--------~~~0----------~

6

7

8

10

18

AGN

Fig. 23 Ranges of stability (in AGN) of phases in alloys of metals in the second and third transition series. From W. Hume-Rothery:J. Less Common Metals, 1964,7, 155.

b

Intermetallic Chemistry of Iron

211

clearly that the bee phase does not tend to go beyond an AGN value of about 8.75. But we have already seen that the fcc phase extends back to an AGN value of about 7.7. Consequently the fields of stability of the fcc and bee phases (in terms of AGN values) overlap and, even if we omit manganese, we no longer have nicely separated fields of stability of the different phases such as we meet in the second and third long periods. 20

Co,wt-CYo 40 50 60

30

70

SO

70

80

90

1400 1200 . Austenite 1100·

u 0

1(XX) -

.: 900 cC.

Ferrite

:::> ~ 800

cC.

UJ

~ 700·

/~~~

" .. ,

UJ

.- 600

"

Ordered 500 400 . 300 200 100 0

0 Fe

Fig. 24

10

20

30

40 50 CO,ot-eyo

60

The system Fe-Co. From: W. Hansen: Constitution of Binary Alloys, McGrawHill, New York, NY, 1958.

We find that if we plot the Fe-Co and Fe-Ni equilibrium diagrams in terms of AGN, the 'Yliquidus curves are accurately, and the 8 liquidus curves roughly superposed, so that there is some sign of a group number effect, but the clear effects of the later periods are not found. This difference between the first and later long periods is probably due to the effect of magnetic contributions to the free energies. In support of this interpretation is the fact that superlattices FeCo (Fig. 24) and FeNi3 exist in the Fe-Co and Pe--Ni systems. Generally speaking, superlattices are formed in systems where the size factors and electrochemical factors are appreciable, and a lowering of energy results if the solute atoms

212

Hatfield Memorial Lectures Vol. II

take up a regular arrangement in which they are as widely separated as possible. Atoms of iron, cobalt and nickel are so similar that the appearance of superlattices is very difficult to understand unless magnetic energy effects are responsible. There is little doubt that in passing from Fe---7Co---7Ni, the fcc structure is becoming increasingly stable, and it is probable that all these fcc structures involve the same kind of hybrid spd orbitals, the proportion of d function decreasing with increasing atomic number. On passing to copper, the fcc structure is retained, but here the bonding electrons are probably of a much more s like nature. As a result of this, although the atomic diameters are entirely favourable, copper and iron form only restricted solid solutions, the equilibrium diagram being as shown in Fig. 25. The very flat liquidus suggests that we are nearing the stage of liquid inuniscibility, and in the system Ag-Fe we find that the two metals are almost completely inuniscible in both liquid and solid states. We know that, in its chemistry, silver is almost exclusively univalent, and the underlying (4d)10 sub-group of electrons is so stable that it cannot be broken into. In contrast to this, copper shows valencies of both 1 and 2, and in the divalent cupric compounds the (3d)10 sub-group is broken down. This suggests clearly that silver does not alloy with iron because it cannot give electrons with a sufficient high proportion of d functions to satisfy the conditions of the iron lattice. In agreement with this interpretation, we find that gold alloys more readily with iron than does copper, and in gold there are well defined trivalent compounds in which two electrons from each Sd sub-group have been removed. Other examples of this kind are known, but these 'electron type' restrictions on alloying can be overcome if two kinds of atom differ appreciably in electronegativity. Fe.wt-% 1600r--------TIo~-=,.20==-----.:::3;.-=-O_40_i_=_--.:5:;..::O~_=;..::60:---7:...=O:..----.,;OO;=_-=OO9=____.

1500 1400 u0l300 •. y

, \ \ \

,

,,

\

,I'

\,1

800 700

~O~~IO~~2~O~3~O--~~~50~+.oo~~m~~8o~~90~~IOO Cu

Fig. 25

Fe,a t -%

Fe

The system Fe-Cu. From: W. Hansen: Constitution of Binary Alloys, McGrawHill, New York, NY, 1958.

Intermetallic Chemistry of Iron

213

Thus, iron and beryllium are able to alloy although there is no possibility of beryllium giving rise to a d type of electron, because no 2d state exists. Here the electropositive beryllium atom attracts the iron atom so strongly that alloying can occur. We may now ask what will happen as we pass on from Cu~Zn~(A1, Ga)~(Si, Ge). The structures of the alloys or iron with these elements serve to show the difference between the behaviour of nickel and cobalt on the one hand, and of iron on the other. We know that, in alloys of copper with many of the B sub-group elements, equilibrium diagrams are formed of which the copper rich parts may be regarded as electron compound diagrams, the positions of the phase boundaries being determined largely by electron concentration with lattice distortion as a complicating factor. Figure 26 shows the equilibrium diagrams of the systems Cu-Zn and Cu-Ga plotted in terms of electron concentration, using the normal valencies of 1, 2 and 3 for copper, zinc and gallium respectively. The almost exact superposition of the diagrams is remarkable, and the ~- and y-phases are typical electron compounds at electron concentrations of 3/2 and 21/13 respectively, while a-solid solution limit occurs at electron concentration 1.4.

-----

1000

Cu-Zn Cu<;O

JJ

~ .. ::>

~0::: ~ ~

QOO a

UJ

t-

Fig. 26

The equilibrium diagrams of Cu-Zn and Cu-Ga alloys plotted in terms of electron concentration. From: W. Hume-Rothery:J. Inst. Met., 1961/1962,9,43.

Figure 27 shows parts of the equilibrium diagrams of the systems Cu-Zn, Ni-Zn, Co-Zn and Fe-Zn. In the systems Ni-Zn and Co-Zn the limits of the fcc solid solutions are at roughly the same atomic percentage of zinc as in the system Cu-Zn, and there are also equiatomic phases with the crystal structures* of typical 3/2 electron compounds. In this part of the diagram nickel and cobalt appear to act as univalent elements, in contrast to the system Fe-Zn where the fcc y-Fe dissolves very little zinc and there is no intermediate phase in the equiatomic region. There is, however, a wide solid solution of zinc in bee aii-Fe; the reason for this is not clear but may be related to that of the Fe-AI alloys referred to below. *~-NiZn has an ordered bee structure. ~'-CoZn has a ~-Mn structure analogous to the 3/2 electron compound in the Ag-Al system. The structure of ~-CoZn is unknown.

214

Hatfield Memorial Lectures VoL II Cu-Zn

1200

a-fcc fJ-bee y-y-bross

800 80 Ni-Zn

100 a. fcc ~. bee ordered CsCI type 13- fcc Cu Au type superlottice "y :a y- brass type

SO

100

a- fcc {3-?

~==~;;:;::::....

800 400

f3 - fJ- Mn

type

..•.•... ~y = 1'- brass

type

0~~~20~~4~0~~601~~~~~

400~

o

I

I 20

I

I

40

I

I

60

(1= 80

I

100

ZINC.,ot-/o

Fig. 27 Cu-Zn, Ni-Zn, Co-Zn and Fe-Zn alloys. From: W. Humc-Rothery: Structure of Alloys of Iron: an Elementarv Introduction, Pergamon, Oxford, 1966.

The

The systems Ni-Zn and Co-Zn contain phases with the characteristic y-brass structure, and it is well known that these fit in with the electron concentration schemes if nickel and cobalt are allotted a zero valency, or at any rate a very low valency. This implies that if nickel and cobalt contribute electrons to the y-brass structure, they absorb a corresponding number into their 3d shells, in agreement with their tendency to complete the (3d) 10 sub-group. In the system Fe-Zn there is a corresponding y-phase but its composition is at a slightly lower percentage of zinc than those of the Ni-Zn and Co-Zn y-phases. This agrees with the smaller tendency of iron atoms to fill their 3d shells. There is no doubt that the equilibrium diagrams of the systems Ni-Zn and Co-Zn resemble that of Cu-Zn much more than does the system Fe-Zn, in agreement with the view that the bonding electrons in iron involve a very high proportion of d orbitals. This conclusion is confirmed by the behaviour of iron, cobalt, nickel and copper with the more electropositive aluminium. The equilibrium diagrams concerned are shown in Fig. 28 and are of great interest.

Intermetallic Chemistry of Iron

215

1200

800

1200

oU•. w

ex:

~ 400~~~~~~~~~~~~ «

ffi ~w

Q..

1600

I-

f200

800

20

40

60

0

ALUMINIUM,ot-/o

80

100

fJ- bee ~/. ordered bee In the system Fe ••AI the bee phose is denoted·ac5

Fig.28

Cu-AI, Ni-AI, Co-AI and Fe-AI alloys. From: W. Hume-Rothery: The Structure of Alloys of Iron: an Elementary Introduction, Pergamon, Oxford, 1966.

In the system Ni-AI the a-solid solubility limit is at almost the same value as in Cu-AI, and that in Co-AI is not very different. Here again, therefore, nickel and cobalt appear to be acting as univalent elements. In contrast to this, the solubility in y-Fe is restricted, just as was the case with zinc. In the systems Ni-AI and Co-AI there are very stable equiatomic phases with ordered bcc structures of the CsCI type. These may be regarded as 3/2 electron compounds of zerovalent nickel and cobalt. At the same time, the ordered CsCI type of structure is one in which each kind of atom is surrounded by eight of the other, and is thus what one would expect for a combination of an electropositive and .electronegative element in which like charges tend to keep away from one another. The great stability. of these

216

Hatfield Memorial Lectures VoL II

phases is the result of the fact that it is only with a trivalent solute that a zerovalent metal gives a 3/2 electron/atom ratio at the equiatomic composition. If we now examine the equilibrium diagram of the system Fe-AI, we see a resemblance to Fe-Zn in that there is a wide solid solution in the bee Fe. This solid solution gives rise to a CsCI type of ordered bee solid solution at the equiatomic ratio, but the structure is much less stable than the corresponding NiAl and CoAl phases; this agrees with the smaller tendency of the iron atom to complete its 3d sub-group, since it is only by absorbing electrons into this sub-group that iron can exert a zero valency which IS required to give the 3/2 electron/atom ratio at the equiatomic composition.

no

Virtually immiscible elements in iron alloys Li Na Mg K Ca

Rb

Sr

Cs

Ba

Ag

Cd

I In

Sn

Hg

Tl

Pb

Bi

The remarkable Fe-AI superlattices are well known and will not be reviewed in detail here. There is the Fe-AI superlattice of the CsCI type referred to above in which aluminium atoms avoid being closest neighbours, and also the Fe3AI type of ordering (Fig. 29) in which both closest and second closest distances of approach are avoided. These structures have been interpreted as the result of aluminium atoms keeping as far away from one another as possible, partly in order to relieve lattice strain and partly to keep the electropositive aluminium atoms away from one another. This is the general explanation of many superlattice structures but we are now led to a very puzzling fact, namely, that in the system Fe-Si (Fig. 30), there is an Fe3Si superlattice similar to but much more stable than that ofFe3AI. This is in spite of the fact that the electronegativities of iron and silicon are very nearly the same and the size factor of silicon relative to iron is very much more favourable than that of aluminium. We may next consider the alloys of iron with the more electronegative elements of Groups IV B-VI B. The elements of Group VI B are all so electronegative that stable compounds are formed at the expense of the solid solutions in ferrite or austenite. Figure

Fig. 29

FeAl and Fe3Al superlattices. From: W. Hume-Rothery: and Alloys, 2nd edn, Cassier, 1948.

Electrons, At01ns, Metals

Intennetallic Chemistry of Iron

217

1600r---------------------------~

OV•. w

a:: 1200

~::::>

? I ~-:Jt)1 ~ '+~t)

< a::

1

w e,

~w t-

-If 020

"

I'CjI I

I

{'

fOOO

II" I

I,

a

" '/

I

.~.., "")

~

41

LI..

I I

'/

800

a'

I I

1030

1/ /1 /1 II" 'I" 'I

a

/,

20

40 SILICON,ot-

Fig. 30

0/0

60

The system Fe-Si. From: W. Hume-Rothery: The Structure of Alloys of Iron: an Elementary Introduction, Pergamon, Oxford, 1966.

31 shows the equilibrium diagram of the system Fe-S, and the pronounced maximum in the liquidus is clear. In the system Fe-O (Fig. 32) even more stable compounds are formed and liquid immiscibility occurs. It is highly probable that molecules ofFe-O exist in the liquid. Ordinary liquid immiscibility such as that in Fe-Ag alloys arisesbecause the two kinds of atom 'do not like each other'. In contrast to this, in a system such as Fe-O, the two kinds of atom attract each other so much that they form a new liquid which does not mix with its parent. In the alloys with elements of Group V B, the equilibrium diagrams of the sequence Fe-P, Fe-As and Pe--Sb show clearly how the maxima in the liquidus become less pronounced with increasing atomic number (Fig. 33). On passing back to Group IV the electronegative nature of the elements is less pronounced, but maxima in the liquidus curves are again found and become less marked on passing from Fe-Si to Fe-Ge. The compound FeSi is of great interest. It has a curious structure in which there is one very short distance of approach, while each atom has six other close neighbours, so that the coordination number is 7. The atomic arrangement of the iron and silicon atoms is the same as that of the Na+ and Cl03- ions in NaCl03. Since the electronegativities or iron

218

Hatfield Memorial Lectures Vol. II

'9130 910°1f----~

a

100

SULPHUR ,ot- 0/0

Fig. 31

The system Fe-S. From: W. Hume-Rothery: The Structure of Alloys of Iron: an Elementary Introduction, Pergamon, Oxford, 1966. Fe-O

SYSTEM

1800~

,

1600

I I I I

I

I

Two

melts

---'\1

1

I I Melt

I~

\Fe

l

6- Fe+Melt 1400~

of.

04

0z

+

02

I

~7-

oU

i::> IZOO~ ~ ~

y- Fe + Wustite

~ 1000~ •....

a i

800~

.,

a - Fe + Wustite

u...

600a-Fe

+

Fe,O.•

400~--~1----~1-----L-1--~-----~1--~-----~1~

o

, Fig. 32

10

20

30

40

50

60

70

The system Fe-O. From: W. Hume-Rothery: The Structure of Alloys of Iron: an Elementary Introduction, Pergamon, Oxford, 1966.

Intermetallic Chemistry of Iron

219

AS,wt-,o

'600r--.IO~~20~~~~

__ ~~ __ ~5~O

6~O~_

.<

N

.,

L&..

u1000

-.

""

T

~ 900 -c

~

c:c

~ 800

~

w

~700

700

I I

a

300~ Fe

6000 10

20

40 30 p.ot-,o

50

60

40

Fe

70

As , at-OJo

so b

Sb,wt-%

0065 70 75 80 85 90 95

-:: 000: I I

800 •

600

,

I

I

I

I'

500 0 Fe

Fig. 33

10

20

so

"

40 50 Sb.or- %

60

70

80

00

100 Sb e

The systems Fe-P, Fe-AI and Fe-Sb. From: W. Hansen: Constitution of Binary Alloys, McGraw-Hill, New York, NY, 1958.

220

Hatfield Memorial Lectures Vol. II

and silicon are almost the same, FeSi is unlikely to be an ionic compound, and its stability is not understood. In alloys of iron, the following metals alloy to such a slight extent that some factor clearly prevents the two kinds of atom from mixing in either the liquid or solid state. If a metal A has a much higher boiling point than a metal B, it has also a higher latent heat of evaporation AA. This indicates strong A-A bonding and so favours the formation of immiscible liquids with an A-rich fraction, unless some factor such as electronegativity difference causes strong A-B bondingt, A markedly unfavourable size factor also favours liquid immiscibility until the difference in atomic diameter becomes so great that interstitialliquid solubility is favoured. Apart from these effects, liquid immiscibility may arise from electronic factors such as those mentioned before in connexion with the system Fe-Ag. The above latent heat and size factor concepts underlie the use of Hildebrand's principle according to which liquid immiscibility is to be expected if VA

+ VB (B _ B )2 > 2RT 2

A

B

Here vA and VB are the molar volumes of A and B, and the solubility parameter BA is defined by BA = (AAlv A)Y2 where AA is the latent heat of evaporation of A. All the above systems agree with this criterion. Of the elements described above the alkalisand alkaline earths are of very unfavourable size factor and of very low A values compared with that of iron. Their behaviour is readily understood, as is that of cadmium and mercury. For thallium, lead and bismuth, the A values are about one half of that for iron, while the size factors are very unfavourable. Further, if they are incompletely ionised, their electrons will be in pure p states and may not be able to combine with the d-type function of iron. The behaviour of silver is probably connected with its exclusively univalent character which may prevent the formation of hybrid orbitals with the proportion of d function required to mix with that of iron; this is in contrast to copper and gold which exhibit higher valencies. In the systems Fe-In and Fe-Sn there are restricted miscibility gaps in the liquid phase for reasons which are not fully understood, but the behaviour of these systems is clearly intermediate between that of the alloys with thallium and lead (almost complete immiscibility) and with gallium and germanium where complete miscibility occurs. We have now completed our survey of the intermetallic chemistry of iron. We are clearly a long way from a satisfactory theory, and there is much need for a more quantitative approach to the subject. Here the clear empirical relations which have been established should be of value to the theoretical workers, and lead them to formulate the real underlying principles. We have reached the stage at which we can begin to survey tThis concept was introduced by B. W. Mott: Phil. Mag., 1957,2,259.

Intermetallic Chemistry of Iron

221

the equilibrium diagrams of iron as a whole and not as a number of isolated individuals. This position has been reached as the result of patient and skilled work by researchers in many countries. It has not been possible to refer to them by name, but many have contributed and this lecture owes much to their work. If Hatfield could visit us today he would be the first to agree that, if it is now possible to give some kind of a sketch, this is because many individuals have worked hard to produce the facts and to try to fit them together. He would admire the view, and would then look down at the map and tell us to get back to work and make it more accurate.

EIGHTEENTH

HATFIELD

MEMORIAL

LECTURE

The Status of the Metallurgy of Cast Irons H. Morrogh At the time the lecture was given Dr Henton Morrogh, F.I.M., F.R.S., was Director of the British Cast Iron Research Association, and Industrial Professor at the University of Technology, Loughborough. The lecture was presented at the Royal Commonwealth Society, London, on 28 November 1967.

I am deeply appreciative of the great honour in being invited to give the Hatfield Memorial Lecture. Other Hatfield lecturers have referred to Dr Hatfield's many contributions to metallurgy, to his great personal attributes, and to his ability to inspire young scientists and technologists. None, however, as far as I can discover, has referred to Dr Hatfield's significant contribution to the metallurgy of cast iron. In 1912 his text book Cast Iron in the Light of Recent Research was published, and for nearly twenty years it remained a standard text book and work of reference. Even today it is a useful starting point in any literature research. Cast Iron in the Light of Recent Research was a great contribution to the rationalisation of a subject originating in those skills and crafts of the ironfounder which made the industrial revolution possible. Dr Hatfield was not only an innovator, he was also a teacher who greatly improved the understanding of the metallurgy of cast irons. For this reason I have chosen the same subject as my theme for this lecture in the hope that I shall not only be able to indicate progress, but also how sound were the foundations of the subject laid by the man whose memory this lecture commemorates.

The tonnage output of iron castings indicates that, next to wrought steels, the cast irons are the most widely used metallic materials of engineering construction. The output of iron castings in the UK has risen from 3 284 000 tons in 1948 to 4 180 000 tons in 1965.1 In the USA during the same period output of iron castings has risen from 14 148 000 to 16 849 212.2 Clearly, the cast irons are of very considerable metallurgical importance, and yet it is probably true that the structure and properties of cast irons are less well understood by the average metallurgist than those of many other materials. In large measure this lack of understanding derives from three features which are peculiar to this family of materials. Firstly, the structure and properties of most cast irons are determined by what happens during solidification, and it is only recently that metallurgists have turned their attention in any systematic manner to the solidification process in metals and alloys. Secondly, the structure of most cast irons is characterised by the presence of graphite

223

224

Hatfield Memorial Lectures VoL II

which (with the possible exception of silicon in, for instance, Al-Si alloys) is almost unique in crystal structure and variety of crystal growth habit among the major phases of the microstructure of the industrial metallic alloys. Finally, the cast irons, belonging essentiallyto the family of high carbon Fe-C alloys, can behave and undergo transformations according to either the stable iron-graphite system or the so-called 'metastable' FeFeC system: the eutectic, for instance, may be austenite-graphite or austenite-cementite. The complexity deriving from this combination of circumstances alone would be unique, but the cast irons usually have, in addition, complex chemical compositions. All cast irons contain five major elements: carbon, silicon, sulphur, manganese and phosphorus, and usually, in commercial materials, many trace elements, either by design or fortuitously, such as lead, tin, antimony and bismuth, and the gaseous elements hydrogen, nitrogen, and oxygen. Many of these minor elements of composition have important and powerful effects. For these reasons the metallurgy of the cast irons often appears mysterious and obscure to the metallurgist who has not had the opportunity to study the field. It seems appropriate, therefore, on this occasion to follow the example set by Dr Hatfield in his book Cast Iron in the Light of Recent Research and to try to penetrate some of the mystery and to show that the behaviour of cast iron, as a result of still more recent research, can be understood in rational terms.

SOLIDIFICATION OF GREY CAST IRONS The cast irons may be very simply defined as the group of high carbon ferrous alloys in which the separation of the Fe-C eutectic takes place during solidification. Although most commercial cast irons are hypoeutectic in composition, with austenite dendrites separating before the eutectic, they may also be hypereutectic, with a separation of primary or 'kish' graphite before the eutectic is reached. The basic features of the solidification process involved in the so called grey cast irons are illustrated schematically in Fig. 1. Superficially this seems very simple, and the general correctness of this sequence is easily established by simple quenching experiments during solidification. Figure 2 shows a cross-section of a 2 in. diameter cylinder of cast iron poured in a sand mould and quenched during the solidification of the eutectic. The light areas are rapidly cooled regions consisting of the cementite eutectic formed during quenching, and the dark regions are the areas of graphite eutectic which formed slowly before the quench. Solidification of the graphite eutectic was more advanced at the edge of the bar than at the centre at the time of the quench. It can easily be seen that eutectic solidification, giving rise to the graphite flakeswhich largely determine the properties of grey cast irons, is nucleated at a discrete number of points, and that growth of the eutectic proceeds on an approximately spheroidal crystallisationfront from each nucleus. This simple description failsto indicate how the typical coarse graphite structures (Fig. 3) formed in slowly cooled castings occur in a form so very much unlike the usual

The Status of the Metallurgy of Cast Irons

Austenite dendrites in liquid

~ ~

cfi

Above the eutectic Austenite dendriteSflj and growing eutectic cells of austenite and graphite;f.

e ~~r: ' (./J /

~.~ ~"'

\

KiS~ gr~phite llquld

1"\

arrest the primary crystals form ••

Kish graphite and growing eutectic cells of austenite and graphite

\

At the eutectic arrest the eutectic of austenite solidifies on a spherical crysta IIi zation front ~raphit~ flakes ~/ In motrix of austenite \

..,,,

-

7."

225

~.:

1

and qraphite

... ~raphite [lokes In an austenite matrix

r

\

On completion of solidification the austeni te of the dendrites and that of eutectic becomes continuous

Fig. 1

Schematic

Fig. 2

representation of the solidification sequences in grey cast iron; left: hypoeutectic; right: hypereutectic.

Microstructure

of 2 in. diameter cylinder quenched during solidification.

eutectic structure. Neither does it indicate how this coarse structure can be converted into very fine eutectiform graphite (Fig. 4) which may be obtained on rapid cooling or by special treatment of the melt.P-" Further, we all know that if we cool a cast iron very rapidly, by casting it into a metal mould, for instance, the eutectic will solidify as the hard,

226

Hatfield Memorial Lectures VoL II

brittle, cementitic eutectic. Further, we know that by nucleating the melt or by increasing the silicon content, it becomes more difficult to cause the cementite eutectic to form. On the other hand, by alloying with chromium, by raising the sulphur content, and by 'denucleating' the melt, it becomes ..easier to form. the carbidiceutectic. We need an explanation of the solidification process which combines and relates these many influences and variables. As a result of research in recent years by many investigators it appears that the broader aspects of a comprehensive explanation are now. available, at least in descriptive terms. We must first consider the 'equilibrium' freezing temperature for the eutectic. In Fig. 5 the equilibrium freezing temperature T'; for the iron-graphite eutectic, and the

Fig. 3

Coarse flake graphite in grey cast iron; etched in picral (xl 00).

Fig. 4 Fine eutectiform graphite in grey cast iron; etched in picral

(Xl 00).

The Status of the Metallurgy of Cast Irons

227

equilibrium freezing temperature T2 for the iron-cementite eutectic, are indicated by two horizontal lines, this arrangement being derived essentially from the 'double' Fe-C; diagram (Fig. 6). The line XU indicates the effect of increasing cooling rate on the Freezing temperature of the iron-graphite eutectic, and WZ the effect of increasing cooling rate on the freezing point of the carbidic eutectic. Extrapolation of XU and WZ to infinitely slow rates of cooling enables the determination of equilibrium temperatures T', and T2 respectively. Figure 5 indicates that with increasing cooling rate the eutectic melt undercools progressively further below Tl for solidification of the iron-graphite eutectic to take place, until at a certain cooling rate indicated by point W, when the melt has undercooled below T2, the solidification of the carbidic eutectic takes place instead of that of the iron-graphite eutectic. As the cooling rate increases and as the amount of undercooling increases from point X to point Y, the number of centres or nuclei from which solidification of the eutectic originates, progressively increases so that in the

~---~------Below thj~ temperature white

iron

~------W~-..-~-r----_LI

can solidih

Y Temperature solidification white iron

COOLING

Fig. 5

RATE

Influence of cooling rate on the temperature Iron.

u

o

240

1220

... Ltqutd

~ I 200 ~ . I 180 ~ 1160

and austenite

Graphite and austenite eutectic may form below this temperature ~

ffi

~

of solidification of the eutectic in cast

1260 1

•...

UJ

of of

1140 IIOO~

3-0

__

~

Fig. 6

__

/ ' /

Carbide eutectic may Liquid_ and form below this temperature carbides ~~~~~~ __ ~ __ ~~ __ ~~ 4·4

Fe-C equivalent double diagram.

228

Hatfield Memorial Lectures VoL II

solidified material the number of eutectic grains or cells, as they are commonly called, also increases. Another feature of the effect of increased cooling rate is that with increased undercooling of the eutectic the growth rate of each eutectic grain or cell also increases. Thus, if we compare a sample of grey cast iron cooled rapidly with another sample from the same melt but cooled slowly, we find that the fonner has relatively few, relatively slowly grown eutectic cells, while the latter will have many more and relatively rapidly grown eutectic cells. Weare now in a position to consider the way the cooling rate or the amount of undercooling influences the structure of the graphite eutectic. Successive stages in the solidification of the eutectic are depicted schematically, with considerable simplification, in Fig. 7. Throughout the solidification process growing graphite ft.akes and austenite are simultaneously in contact with the liquid at the periphery of the growing eutectic cells. This growth process can be shown, by simple quenching experiments, to apply to irons with coarse ft.ake graphite and to irons with fine ft.ake graphite. Figure 8 shows a eutectic cell with fine graphite in a sample quenched during solidification. Russian workers> were the first to emphasise that within each eutectic cell there is a continuous skeleton of graphite of the form shown in Fig. 9, which repeatedly branches or otherwise divides during the growth of the eutectic cell. With slow rates of growth, that is generally under conditions of little undercooling and slow cooling rate, the frequency of branching is relatively infrequent and, in the microstructure, coarse ft.akes, some of which will appear disconnected, will result. With considerable undercooling and high rates of growth arising from rapid cooling, the branching process is much more frequent and fine graphite structures result. The interconnected character of the graphite in each eutectic cell has been shown by different techniques by several workcrs.f=? Figure 10, taken with a scanning electron microscope, shows the graphite skeleton in a eutectic cell with coarse graphite flakes. This picture was taken by first etching away the metallic matrix. Figures 11 and 12 show the interconnected branched arrangement in a eutectic cell with very fine graphite revealed by the same technique. Since it is usually the object of the foundryman to produce machinable castings, free from the carbidic eutectic, he will try to arrange events so that point W (Fig. 5) occurs at the highest cooling rate possible, taking into account other factors such as the need to adjust the chemical composition of the iron to suit specifications for mechanical

a

Fig. 7

b

c

Successive stages in the solidification of the flake graphite eutectic in cast iron.

The Status of the Metallurgy of Cast Irons

Fig. 8

229

Cast iron with fine eutectifonn graphite quenched during solidification of eutectic.

Fig. 9

Drawing of graphite skeleton in a eutectic cell.

properties. We are now able to provide a rationalisation in descriptive terms of the various features that the ironfounding metallurgist controls to meet this requirement. Firstly, from Fig. 5, it is obvious that the foundryman stands a better chance of avoiding the formation of the carbidic eutectic if he can either reduce the amount of undercooling for a given cooling rate, or ifhe can arrange that the temperature difference between the equilibrium freezing temperature for the iron-graphite eutectic T1, and that of the carbidic eutectic T2, is as large as possible. If the temperature interval is large, the eutectic liquid can undercool more, that is can sustain higher cooling rates before it cools into the region of white iron solidification. Although he did not think in these terms, the foundryman of fifty years ago, thanks to the pioneer researches of Professor Turner, 8

230

Hatfield Memorial Lectures VoL II

Fig. 10

Graphite skeleton in coarse graphite iron after removal of metal; taken with scanning electron microscope (x260).

Fig. 11

Graphite skeleton in very fine graphite iron after removal of metal; taken with scanning electron microscope (x 500).

knew how to achieve the desired result by appropriately raising the silicon content of the iron. As is shown in Fig. 13, with increasing silicon content the temperature range over which the iron-graphite eutectic can freeze before undercooling into the region of carbidic eutectic solidification increases with increasing silicon content.

The Status of the Metallurgy of Cast Irons

Fig. 12

231

Same as Fig. 11, but at a higher magnification (x 2300).

1160

u o

Graphite austenite eutectic equilibrium temperature

0-5

Fig. 13

1-0

1-5 51LI CON,

0/0

2·0

2·5

Effect of silicon on eutectic equilibrium freeing temperatures.

Some elements, such as copper.!? have an effect similar to that of silicon, while other elements, such as chromium, have an opposite effect,"! reducing the temperature interval between the two eutectics, as shown in Fig. 14. To prevent the formation of the unwanted carbidic eutectic in grey cast irons, the foundryman tries generally to avoid the presence of elements such as chromium, and maintains the silicon content at a fairly high level. Unfortunately, there is a limit to which silicon can be used in this way, since beyond certain levels it has unwanted effects, particularly on mechanical properties. Thus, to prevent the formation of eutectic carbides, particularly in high strength irons which must have relatively low silicon contents, the ironfounder needs to use other techniques, which generally involve the nucleation of the melt by a process known as inoculation.

232

Hatfield Memorial Lectures VoL II

This usually involves the addition of a small amount of a silicon containing material such as ferrosilicon or calcium silicide or graphite to the melt shortly before casting. Such additions increase the number of nuclei available for the solidification of the graphite eutectic, and as a result of such additions there is an increase in the number of eutectic cells visible in the microstructure. Figure 1512 shows the effect of increasing additions of 75-80% silicon containing ferrosilicon, expressed as percentage silicon increment, on eutectic cell number and hence on nucleation. Such increases in nucleation reduce the amount of undercooling of the eutectic liquid before solidification of the graphite eutectic begins at a given cooling rate. This effect is depicted schematically in Fig. 16, in which iron 1 represents an iron which has not been inoculated, and iron 2 represents inoculated material. At a given cooling rate, say R3, the well nucleated iron 2 undercools much less before solidification of the eutectic than iron 1, and as a result can sustain a very much higher cooling rate R2 before undercooling into the region of carbidic solidification. The production of iron castings involves compromise. To avoid eutectic carbide a relatively high silicon content is required. To produce a high tensile strength a relatively low silicon content is required. To avoid eutectic carbides with low silicon contents,

0-4

0-6

CHROMIUM,

Fig. 14

io

0-8

0/0

1-2

Effect of chromium on eutectic equilibrium freezing temperatures. cI2 1--

Z

~ 0-5 ~ 0-4

u

~ 0-3

3900

4 800

5 700 EUTECTIC

Fig. 15

6 600 CELLS

7 500

/ in2

Effect of additions of ferrosilicon to the molten iron before casting on the number of eutectic cells in cast iron.

The Status of the Metallurgy of Cast Irons

233

LU

ex:: :::> I-

«

ex::

R4

R3

LLI Q..

::E LLI

I-

Rz

R, COOLING

Fig. 16

RATE

Showing effect of increased nucleation (iron 2) on eutectic solidification temperature with various rates of cooling.

eutectic nucleation may be increased by inoculation, but again the foundryman cannot always make unrestricted use even of this technique, for other very important practical reasons. Grey iron castings tend to swell during solidification in a sand mould. The amount of swelling is less, the more rigid the mould. If the mould lacks sufficient rigidity, expansion of the casting during solidification can cause unsoundness in the form of internal porosity or surface sinking defects, particularly in those parts of a casting last to solidify. 13· The fonnation of a simple cavity by this process is illustrated diagrammatically in Fig. 17. For a mould of a given rigidity the amount of swelling, and hence the probability of developing unsoundness in the manner shown in Fig. 17, increases with the number of centres of eutectic solidification.l+ Figure 18 shows the effect of increased eutectic nucleation on the diameter of spherical test castings poured in irons of various compositions. The inoculated spheres show significantly larger dimensions than the spheres made from metal not inoculated, and hence of relatively low nucleation. It should not be assumed that it is impossible to produce sound castings from highly nucleated metal. Everyday experience shows that such an assumption would be wrong. It merely becomes more difficult to produce sound castings if the degree of nucleation is high. Nevertheless, sometimes this effect manifests itself in a quite dramatic manner, as is shown by the two pictures. reproduced in Fig. 19. These show sections through a commercial clutch plate casting. The casting on the left has been inoculated, has a relatively large number of eutectic cells, and has considerable internal unsoundness. The casting on the right was made from metal for which the inoculation procedure was omitted; it has fewer eutectic cells and is completely sound. This effect of the number of centres of eutectic solidification on the tendency to swell during solidification and the resulting tendency to give internal unsoundness, appears to operate however the nucleation of the graphite eutectic is varied. The number of centres of eutectic solidification may be increased by additions of nucleation agents, but it can also be increased by changes in chemical composition. An increase in sulphur contentlv-"? increases the number of centres of eutectic solidification, for instance, and this in tum

234

Hatfield Memorial Lectures Vol. II PLAN

B:•••

ELEVATION

( b)

(0)

(c)

Fig. 17 Formation of internal cavity in grey cast iron due to expansion of casting during solidification. (a) mould filled with liquid metal immediately after casting; (b) solidification commenced from mould walls inwards; (c) hot liquid metal flows from the boss to compensate for mould wall movement, producing a shrinkage cavity; (d) casting completely solidified with casting remaining in boss. 3·040 .s

•

,-~----

x

__ ~--

"..---x--,-

e __ ----

.,-

UJ

0:: u.J

~ 3·000 V'l

4·00 CARBON

Fig. 19

---;(--l(-

Inoculated

-e- --e- Uninoculoted 3-90

Fig. 18

•

4·10 EOUIVALENT,

4·20

0/0

4·30

Effect of nucleation on dimension of grey iron castings.

Effect of nucleation on soundness of clutch plate castings. (a) inoculated; (b) untreated.

The Status of the Metallurgy of Cast Irons

235

increases the tendency to unsoundness.I? Conversely, the number of centres of eutectic solidification, that is the degree of nucleation, may be deliberately reduced by violent agitation of the melt-? before cooling or by the addition to the melt of a small amount of titanium.3,20 All these treatments tend to give sounder castings, but by reducing the nucleation of the melt they promote increased undercooling before solidification of the eutectic and hence increase the possibility of undercooling into the region of carbidic solidification. In practical terms this means that, while these techniques of reducing the degree of nucleation may promote sounder castings, their use makes it more difficult to produce castings free from the carbide eutectic. The foundryman therefore requires a method of treating his metal so that it solidifies with as little undercooling as possible (this gives the greatest possibility of avoiding the carbide eutectic) and so that solidification of the graphite eutectic takes place from as few centres as possible (this gives the best chance of obtaining a sound casting). Some progress has been made in this direction, but before this can be described it is necessary to take account of another factor influencing the tendency of the eutectic liquid to undercool into the region of carbidic solidification. We have already seen (Fig. 5) that fast cooling encourages undercooling into the metastable region, that this possibility is made more difficult by adjusting the composition of the melt to give as big a temperature difference as possible between the stable and the metastable eutectic freezing temperatures (Fig. 13), and that nucleation of the melt also reduces the chances of this happening (Fig. 16). The amount of underco oling for a given rate of cooling and a given degree of nucleation must also be influenced by the growth rate of the eutectic cells. If the growth rate is increased, solidification takes place with less undercooling, and, if the growth rates of the eutectic cells are decreased, the amount of undercooling will tend to increase. The evidence available= suggests that increases in sulphur and hydrogen, for instance, may reduce the growth rates of the eutectic cells of the graphite eutectic at a given degree of undercooling. This would explain the marked tendency of both of these elements to encourage the formation of the carbide eutectic and to promote the formation of coarse flake graphite structures20,21 where otherwise fine flake graphite structures would be expected. It now appears that we are in a position to take practical advantage of this growth rate effect. Little or nothing is known of the nature of the nuclei which are responsible for the nucleation of the iron-graphite eutectic. The effects of superheating in eliminating nuclei have been held by many workers22-24 to be the result of the elimination of small particles of oxides and silicates from the melt. It is perhaps not surprising that graphite additions themselves can function as nucleating agents, but the most commonly used nucleating additions are usually based on ferrosilicon having 60-80%Si. For ferrosilicon to function effectively in this way it must contain a small amount of aluminium and calcium.s'' and most commercial ferrosilicon inoculants have a formulation of this type. In the last few years ferrosilicon inoculants containing barium have been developed, and more recently it has been found that ferrosilicon containing strontium has a powerful nucleating effect.26 This powerful influence of the strontium is illustrated in Fig. 20, which shows sections of

236

Hatfield Memorial Lectures Vol. II

the edges of two o/t6in thick plates, the top plate being nucleated with 0.1 % normal calcium and aluminium containing ferrosilicon, and the bottom plate with the same amount of strontium containing ferrosilicon. The white areas are the carbidic eutectic formed due to the rapid cooling at the edge of this very thin section. It can be seen that the plate cast from metal treated with the strontium ferrosilicon has many more eutectic cells and much less of the carbide eutectic than the plate treated with the normal ferrosilicon.

Fig. 20 inoculated

Carbidic eutectuc formation in 0/16 in plates; etched in Stead's reagent (x5). top: with 0.1 % normal ferrosilicon; bottom: inoculated with 0.1% strontium ferrosilicon.

More recently it has been shown-'? that the tendency for the carbidic eutectic to fonn can be reduced by nucleating additions of strontium ferrosilicon, even though the number of eutectic cells growing may be the same as or less than those resulting from a nucleating addition of normal ferrosilicon. This is shown in Fig. 21, in which it can be seen that, in this instance, the amount of the carbidic eutectic is similar and the number of eutectic cells fewer in the sample nucleated with strontium ferrosilicon than in the sample treated with a normal inoculant. If the tendency to unsoundness is determined by the number of eutectic cells growing, we have here an opportunity to reduce the tendency to unsoundness and at the same time reduce the tendency for formation of the carbidic eutectic. That this can in fact be achieved is shown in Fig 22, in which the tendency to give the carbidic eutectic, expressed as 'depth of chill', is plotted against the tendency to unsoundness, expressed as 'depth of sinking' for two series of test samples, one of which had various amounts of normal ferrosilicon and the other various amounts of strontium ferrosilicon as nucleating agents. It can be seen that for the same proneness to give the carbidic eutectic, the severity of unsoundness is much less with the strontium than with the normal ferrosilicon. The mechanism of the influence of the strontium inoculant may be deduced from the data given in Table 1. Examination of these data reveals that at each cooling rate the number of nuclei growing is less and the amount of undercooling is less for the strontium

The Status of the Metallurgy of Cast Irons

Fig. 21

237

Similar carbide eutectic amount and fewer eutectic cells with strontium ferrosilicon inoculant. (a) strontium ferrosilicon; (b) normal inoculant.

3-500

4-00

in x 102

Fig. 22

Showing reduced tendency to eutectic carbide formation and reduced tendency to unsoundness with strontium ferrosilicon as nucleation agent.

238

Hatfield Memorial Lectures VoL II

ferrosilicon than for the normal ferrosilicon. This strongly suggests that the strontium is having an effect on the growth process, and clearly, from the information presented earlier, the material treated with the strontium inoculant would be less prone to unsoundness due to the lower number of nuclei growing, and less prone to enter the temperature range of metastable solidification because of the reduced undercooling. Table 1

Effect of additions ofO.2%Si as normal ferrosilicon and as strontium ferrosilicon on number of eutectic cells and undercooling at various cooling rates

Cooling rate, degC/min

Normal ferrosilicon no. of undercooling eutectic cells °C

50 100 300 400

750 1200 5800 6800

11.5 13 27.5 34.5

Strontium ferrosilicon no. of undercooling, eutectic cells °C 550 1000 5200 6300

10 12 25.5 32

MECHANICAL PROPERTIES OF GREY CAST IRONS So far we have seen how the essential characteristics of ordinary grey cast irons are determined during the solidification process by the interplay of the effects of composition, cooling rate, nucleation and eutectic growth rate, and the possibility of solidification of the material according to the stable or the metastable system. Additionally, it has been shown that after solidification the basic feature of the microstructure of grey cast iron involves extended branched skeletons of graphite locked together usually by dendrites, as shown in Fig. 23. This array confers somewhat unusual behaviour on cast irons when stressed.

Fig. 23

Drawing of a graphite skeleton in a eutectic cell with interpenetrating primary dendrites.

The Status of the Metallurgy of Cast Irons

239

The stress-straincurve for all flake graphite irons appearscurved from the origin, with no straight line portion indicating purely elastic behaviour. By stressingto and understressing from progressivelyincreasing stresses,and measuring each time the strain at the maximum stress and the strain after removing that stress,it is possible to show that the total strain is composed of permanent deformation and recoverable deformation, as shown in Fig. 24, for a sample tested under tensile loading. Longitudinal and lateral stress-strain curves are indicated. The recoverable strain curve in the lateral direction follows a straight line and may be assumed, therefore, to represent elastic deformation of the metallic matrix. However, the recoverable strain curve in the longitudinal direction is not a straight line, and therefore a more complicated process than simple elasticdeformation is involved.

0·4

Fig. 24

STRAIN,

0-6

010

0

0-1

0·2

Permanent and recoverable components of tensile strain for a grey cast iron.

It has,been shown28-3o that grey cast irons undergo permanent volume changes under stress and that these volume changes are associated with changes in shape of the spaces occupied by the graphite eutectic skeletons. In the longitudinal direction, under tensile stress the spaces occupied by the graphite increase in volume, while there is little or no change in volume in the lateral direction. Under compresive stress there is little or no change in volume in the longitudinal direction, but an increase in the space occupied by the graphite in the lateral direction. Using this concept, Gilbert.e? by measuring total, permanent and recoverable' strains in the longitudinal and lateral directions under tensile and compressive loading has been able to show that deformation in the longitudinal direction under tensile stress and in the lateral direction under compressive stress may be resolved into four components of strain, as shown in Figs. 25 and 26. In these directions we have normal elastic and normal plastic deformation, but in addition there is recoverable and permanent strain resulting from ·the change in volume during stressing of the spaces occupied by the eutectic graphite skeletons. The mechanical properties of grey cast irons may be varied by alloying with various elements in a manner similar to that applying in other ferrous materials, particularly to control or modify the structure of the metallic matrix, but the graphite structure has an overriding determining influence on the properties of the material. The graphite

240

Hatfield Memorial Lectures VoL II

V") V") UJ

a::: l-

V")

uJ .-J V')

~I- 2 0·4 STRAIN,

Fig. 25

Components

0·6

0/0

0

0·2

of tensile strain associated with the matrix and with the volume of the spaces occupied by graphite.

lV')

uJ

>

(/) (/)

uJ

a::: 0-

L

o

U

Fig. 26

Components

1·2 0 STRAI N,

0-4

0/0

O-S

1·2

of compressive strain associated with the matrix volume of the spaces occupied by graphite.

and with the

structure derives from the solidification process. The extent to which alloying elements may be deployed to modify structure and properties is limited to their likely effects on the solidification process. Furthermore, elements of composition which in other ferrous materials may be unimportant or have only limited influence, may modify profoundly the structure and properties of cast iron by their influence on the solidification process and on the growth of graphite. The most obvious instance of this is the production of irons with spherulite graphite by solution of magnesium, sometimes together with rare earth elements, in the molten iron. By this procedure the crystal growth habit of the graphite is substantially modified. Normally, graphite grows by crystallisation from the melt in the direction of the close packed basal planes. In a graphite spherulite (Fig. 27) there is a radial array of twisted fibres of graphite growing apparently from a common centre and which may have the structure shown schematically in Fig. 28. Each graphite fibre appears, at least superficially,

The Status of the Metallurgy of Cast Irons

241

to be growing at right angles to the normal growth direction, that is at right angles to the basal plane.>! There are many theories attempting to explain how the presence of elements such as magnesium have this effect. It is not possible to go into this subjecthere, but the explanations offered include the possibilities of the occlusion of foreign atoms in the growing graphite, the creation of special nuclei or nucleating conditions, and the provision or removal of surface active agents at the surface of the growing graphite crystals. None is yet completely satisfactory and much more experiment is required. For the purpose of this lecture I want simply to emphasise that the effect is produced by a very small amount of the element concerned (0.02-0.04% in the case of magnesium). This small addition to the composition, and the accompanying composition changes necessary for the solution of magnesium (reduction of sulphur (and possibly oxygen) to very low levels) changes the growth habit of the graphite, profoundly influences the mechanical properties such that the resulting cast irons behave as more or less normal ductile ferrous materials, and completely changes the character of the 'eutectic' solidification process. In irons with spherulite graphite, the number of centres of eutectic crystallisation is usually greater by a factor of about 200 than for corresponding flake graphite cast irons,32,33 and furthermore the growth of the eutectic takes place by a completely different mechanism. Whereas in the flake graphite irons growth of each eutectic cell takes place with graphite austenite and liquid in simultaneous mutual contact, as shown in Fig. 7, in the spherulitic graphite irons the growing graphite spherulites quickly become surrounded by shells of austenite, as shown in Fig. 29, and solidification of the eutectic proceeds by the transport of carbon from the liquid across the austenite shell to the graphite spherulite. This is a relatively slow process and, as a result, in spite of the large number of nuclei growing, the eutectic liquid can undercool very easily into the region of metastable solidification.

Fig. 27

Graphite spherulite (xlS00).

242

Hatfield Memorial Lectures VoL II

Fig. 28

Schematic depiction of probable array of twisted fibres in a graphite spherulite.

a

Fig. 29

b

c

Apparent solidification process for the eutectic with spherulitic graphite.

The irons with spherulitic graphite have demonstrated the extraordinary sensitivity of the crystallisationof graphite to the influence of traces of impurities. The presence of traces of elements such as lead, antimony and bismuth>+can prevent the production of graphite spherulites by magnesium, giving flake graphite structures instead. However, the harmful effectsof these elements can be neutralised by the addition of a small amount of a rare earth element along with the magnesium, but, if no hannful elements such as lead, antimony or bismuth are present, the rare earth elements can inhibit the development of good graphite spherulites, particularly under conditions of slow cooling. The situation becomes even more confusing and complex with the finding that well fonned spherulites and a high degree of nucleation can best be obtained by the addition of bismuth, provided a rare earth element is present. These are not imaginings based on imperfect experiment, but are well established observations of which considerable use is made in the industrial production of spherulitic graphite cast irons.P> The sensitivity of the solidification process in cast irons and the susceptibility of the graphite phase to modification by traces of impurities is not limited to irons with

The Status of the Metallurgy of Cast Irons

243

spherulitic graphite, but also applies to irons with flake graphite. The solidification of the flake graphite is powerfully influenced, for instance, by nitrogen content.37,38 Figure 30 illustrates this effect of nitrogen, which causes the individual graphite flake to become shorter and thicker (apparently growth in the direction of the basal plane is progressively discouraged and growth at right angles to this becomes more pronounced). This effect of nitrogen can have important effects on the mechanical properties of grey cast iron and, for instance, an increase in nitrogen from a normal value of about 0.005 to 0.015% can give an increase in tensile strength of as much as 30%.

Fig. 30

Effect of nitrogen on shape of graphite flakes; etched in picral (x300). (a) O.008%N; (b) O.015%N; (c) O.03%N.

The separation of the graphite phase during solidification can also be influenced by traces of impurity with adverse effects on mechanical properties. If lead in amounts in excess of 0.003% is present in ordinary flake graphite cast irons, and if the hydrogen content at the same time is higher than the normal 0.0001-0.0002%, some of the graphite separates in a pseudo- Widmanstatten pattern, as illustrated in Fig. 31. This regular array of graphite can cause a 50% drop in tensile strength. The presence of lead has very important practical significance, on which considerable experimental work has been carried out, but the mechanism by which this element gives rise to the structure is not known at present. When the structure was first observed it was assumed to be a true Widmanstarten structure, but evidence is accumulating that it may be a regular separation of the graphite eutectic.

CONCLUDING REMARKS In this lecture I have tried to emphasise that, superimposed on the normal alloying effects of the elements commonly present in cast iron and which have been understood in general terms at least for a long time and were well established in Dr Hatfield's book, there are a number of other phenomena which have very important effects, and which

244

Hatfield Memorial Lectures VoL II

Fig. 31

Regular array of graphite due to presence of lead; etched in picral (x600).

concern particularly the metallurgist dealing with cast irons. These phenomena derive in large measure from the solidification of the graphite eutectic, and the special nature of graphite itself Unfortunately the systematic study of solidification of metals and alloys has not been particularly fashionable until the last few years. Suddenly, however, widespread interest has developed and I have taken the opportunity to draw attention to the cast irons as important engineering materials and upon which much further fundamental work on the solidification process could usefully be carried out. Cast irons are not mysterious, uncontrolled materials and, when the solidification process is properly appreciated, the cast irons behave in a fairly normal way, and the cast iron metallurgist has a firmly established and rationally organised corpus of knowledge enabling the control of the materials, sometimes in quite subtle ways. The cast irons are remarkably sensitive to the effect of trace elements on the solidification process and the separation of graphite. The practical application of this knowledge is leading by a long way the scientific understanding of some of the phenomena observed. This will be a rewarding field for further research. At the time of the appearance of Dr Hatfield's book Cast Iron in the Light oj Recent Research the salient features of the effects of the common and alloying elements in cast iron had been established. In general outline, the pattern then stated applies today. Since then, however, progress has been particularly in the direction of obtaining a better insight into the materials by a better understanding of the solidification process and the effects of traces of impurities on this. This is the aspect which I have tried to emphasise. To venture into these fields is possible today due to the sound foundations laid over half a century ago by metallurgists of whom Dr Hatfield is an outstanding example.

The Status of the Metallurgy of Cast Irons 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. 33. 34. 35. 36. 37. 38.

P. MOORE: Brit. Found., June 1966, 59, 254-282. Marketing Guide to the Metal Casting Industry; Penton, Cleveland, OH, 1967. A. L. NORBURY and E. MORGAN:JISI, 1936, 134, 327P-346P. J. V. DAWSON: BClRAJ. Res. Dev., June 1956,6,249-258. K. P. BUNIN et al.: Lit. Proiz., 1953,4, (9),25. W. OLDFIELD:BCIRAJ., March 1960, 8,177-192. M. G. DAY: unpublished work. T. TURNER:J. Chern. Soc., 1886,99,215. W. OLDFIELD:BClRAJ., Jan. 1962,10,17-27. K. LOHBERG and K. LOHRIG: Giess. Techn. Wiss., 1966, 18,63-82. W. OLDFIELD:BCIRAJ., July 1961, 9,506-518. A. G. FULLER:BCIRAJ. Res. Dev., Feb. 1958,7,157. J. M. GREENHILL:BCIRAJ., March 1962,10,158-165. K. E. L. NICHOLAS:BClRAJ., March 1962, 10,166-172. J. D. BERRY et al.: BClRAJ., May 1960, 8, 377-392. A. G. FULLER:BClRAJ. Res. Dev., Feb. 1958,7,157-170. W. OLDFIELD:BCIRAJ., March 1960, 8,177-192. H. MORROGH: Brit. Found., May 1960,53,221-242. A. G. FULLER: BCIRAJ. Res. Dev., June 1959, 7, 725-733. A. BOYLES: The Structureoj Cast Iron; ASM Cleveland, OH, 1947. W.J. WILLIAMS:JISI, April 1950, 164,407-422. O. VON KEIL et al.: Arch. Eisenh., 1933-34,7, 579. K. LOHBERG et al.: Giess. Techn. Wiss., Jan. 1964, 16, 15-34. A. DE Sv: Mod. Castings,July 1967, 67-78. J. V. DAWSON: BCIRAJ., March 1961, 9,199-236. J. V. DAWSON: Mod. Castings, May 1966, 49,171. R. A. CLARK and T. MCCLUHAN: Mod. Castings, 50, Sept., 394-400. L. F. COFFIN:J. Appl. Mech., Sept. 1950,17,233-248. W. R. CLOUGH and M. E. SHANK: Trans. ASM, 1957,49,241-262. G.]. N. GILBERT: BCIRAJ., July 1963,11,512-524. H. MORROGH: BClRAJ. Res. Dev., June 1955, 5, 655-673. I. C. H. HUGHES: Proc.Inst. Brit. Found., 1952,45, A.157-A.174. W. OLDFIELDandJ. G. HUMPHREYS:BClRAJ., May 1962,10,315-324. H. MORROGH: Trans. AFS, 1952,60,439-451. C. R. LOPER and R. W. HEINE: US Patent No.3311469, March 28, 1967. J. V. DAWSON et al.: BClRAJ. Res. Dev., June 1953, 4, 540-552. F. A. MOUNTFORD: Brit. Found., April 1966, 59,141-151. I. C. H. HUGHES and G. HARRISON: BClRAJ., May 1964,12,340-360.

245

TWENTIETH

HATFIELD

MEMORIAL

LECTURE

Metallic Chemistry in One, Two and Three Dimensions L.

s. Darken

At the time the lecture was given, Dr Darken was with the Edgar C. Bain Laboratory for Fundamental Research, United States Steel Corporation Research Centre, Monroeville, Pennsylvania, USA. The lecture was presented at Sheffield University on 24 November 1969.

It is a great pleasure, honour and distinction to be privileged to deliver this 20th Hatfield Memorial Lecture. Although, chronologically speaking, it would have been possible, I never actually had the opportunity to meet Dr William H. Hatfield, though his fame was widespread at the time I started my metallurgical career in 1935. In retrospect, it is truly astonishing to us or at least to me, in this latter day, that one man could have achieved so much, writing well over 100 publications in a wide range of metallurgical areas and later chairing the joint research committees and actively participating in the preparation of their famous Special Reports of the Iron and Steel Institute, all while developing and directing the Brown-Firth Research Laboratories here in Sheffield. Professor Waterhouse, in the first Hatfield Memorial in 1946, stated that Dr Hatfield, during the war, 'was largely responsible for convincing the authorities that, in spite of the concentration of effort on increased iron and steel production, it was necessary for research work to continue.' Such persuasive power is not the least of the many attributes for which we miss Dr Hatfield today. I believe he would have lent his support to the sort of investigations I shall discusshere, for he realised that we must ever be in the process of building for the metallurgical future. As most of you probably guessed, when you heard I was to speak here on a topic of my own choosing, that topic would be involved with thermodynamics, no matter what the title might be. Such is indeed the case, and my focus will be first on the thermodynamics of solutions of various types, especially those involving 'anomalous' behaviour, and second on the bearing of such thermodynamics on the kinetics of some metallurgical reactions.

A few years ago, when he was 80 years old, Joel H. Hildebrand! began a famous lecture with an exposition of the aesthetic aspects of science. Hildebrand is one of our great physical chemists and educators, best known to some as the author of the book, 'Solubility of non-electrolytes'. In this lecture he briefly reviewed a major proportion of his life's research by showing a plot of the solubility of iodine, expressed as the logarithm of 247

248

Hatfield Memorial Lectures VoL II

its mole fraction, commented:

against the square of the difference

of solubility

parameters.

He

The points. . . represent the solubility of iodine in fifteen solvents in which it dissolves with a beautiful violet colour. The line through them represents the theory for such solutions, developed when only a few of the points had been determined . . . This plot, here made public for the first time, is an example of 'the profounder beauty which comes from the harmonious order of the parts'. This was Henri Poincare's expression: ' ... the profounder beauty which comes from the harmonious order of the parts.' I hope that besides utility, a trace of such beauty may be found in some of the examples to be discussed, and that you will forgive me if this lecture sounds like the review it dominantly is, of some recent work of our laboratory. The title of my lecture was obviously chosen well in advance of its preparation and was intended to cover any of the multitude of sins I might subsequently choose to commit today in the micro century allotted to me. In addition to possessing this virtue, my title was also intended to emphasise the fact that metallic chemistry is importantly involved not only with (a) our three-dimensional solutions composed of essentially zerodimensional atoms but also with (b) the essentially two-dimensional surfaces and interfaces where heterogeneous reactions occur, and further with (c) the dislocations, which, with their stress fields may, perhaps less accurately, be described as one-dimensional. To illustrate some of the features to be involved, I have taken the liberty of borrowing some illustrations from a text by Matteoli- and of adding a few in the same style (Fig. 1).

OBSERVATIONS BY FIELD ION MICROSCOPY Many of these features, as I'm sure you are aware, may now be observed 'directly' by field ion microscopy, which reveals the sites of individual atoms on a finely pointed tip as in Fig. 2, from Brenner. 3 A grain boundary (indicated by arrow) in a tungsten alloy is seen in the field ion micrograph-' of Fig. 3. A typical feature is the narrow 'width' of the boundary: hardly more than the diameter of a single atom. Iron is more difficult to image since the imaging voltage is close to the voltage for field evaporation. However, with the aid of recently developed channel plate image converters, lower imaging voltages can be used yielding improved field ion micrographs of iron and other less refractory metals; a grain boundary in iron, as revealed by this technique? is shown in Fig. 4. An emerging screw dislocation in a low alloy steel is shown> by the spiral terraces of Fig. 5. By appropriate modification of the field ion microscope the image of a selected atom may be moved to fall on a small hole in the fluorescent screen. Upon voltage pulsing, the corresponding atom flies through the hole and into a mass spectroscope by which its chemical identity may be established. Such remarkable identification was first achieved by Muller et al.6 and more recently in an apparatus of modified design by Brenner and Mclvinney.? The atom by atom analysis has already yielded some surprising results. After

Metallic Chemistry in One, Two and Three Dimensions

c

a

Fig. 1

Features of crystalline and other states (from Matteoli? and in same style).

Fig. 2

Field ion micrograph: Tungsten (from Brenner").

249

250

Hatfield Memorial Lectures VoL II

Fig.3

Fig.4

Grain boundary in W-3%Re alloy (from Brenner").

Neon image of iron using channel plate (from Brenner and Mcls.inney-').

Metallic Chemistry in One, Two and Three Dimensions

Fig. 5

251

Screw dislocation in a low alloy steel; H2 image (from Brenner and Perepezko"); micrographs taken sequentially after successful field evaporations.

nitrogen exposure, a tungsten surface shows many randomly located atoms, such as shown in Fig. 6. Brenner and Mclvinney+ found that nearly all of these atoms are displaced substrate atoms and not nitrogen atoms as had previously been assumed. The atom probe is not restricted to the analysis of adsorbed atoms but can be used to measure chemical fluctuations much in the manner of the well known electron microprobe. But now the spatial resolution is angstroms instead of microns and the detection of segregated impurities at grain boundaries becomes possible as shown by the two spectra in Fig. 7, taken of a phosphorus-containing steel near a grain boundary and away from the boundary.f

LIQUID AND SUBSTITUTIONAL SOLID SOLUTIONS Any discussion of the thermodynamics of metallic solutions can hardly avoid a treatment in terms of the activity of the components. The activity, though defined in terms of the chemical potential or partial molal free energy, is easily visualised as proportional to the

vapour pressure and is frequently measured by a vapour pressure technique.

252

Hatfield Memorial Lectures VoL II

Fig.6 Field on micrographs of emission centres on W(110) after N2 condensation; atoms indicated by arrows were revealed by atom probe to be W (from Brenner and Mck.inney+). GRAIN BOUNDARY 56

1-5

2-0

2-5

~+2at-oloCr+0-07at~'('ALLOY

Fig. 7

3-0

t,J.ls

1-5

2~0

ANNEALED 250h

3-0 52SoC

Atom probe spectrum ofFe-2 at.-%Cr+0.07 at.-%P alloy, annealed 250 h, 525°C, showing segregation ofP at grain boundary (from Goodman et al.8).

The standard state is frequently chosen as the pure substance, and for an ideal liquid or substitutional solid solution, the activity of any component is then equal to its atom fraction; the plot of activity v. atom fraction would then be a beautiful straight line, known as the Raoult law line. Such beauty, alas is rare, as indicated by the plots from recent data9-16 in Fig. 8, for liquid iron solutions. However, the proportionality of activity to concentration (Henry's law) at the high iron side and the approach to the Raoult law line at the low iron side are quite evident, though the linear ranges are indeed short.

Metallic Chemistry in One, Two and Three Dimensions

253

x

o

>-•.

.... 5>

5«

Fig. 8

Activity: liquid iron alloys.9-16

Some semblance of order is introduced by invention of the activity coefficient defined as the ratio of the activity to atom fraction. In fact, we have recently shown 17-19 that, over a substantial range of compositions near each terminus, the logarithm of the activity coefficient of one component of a binary liquid or substitutional metallic solution is a linear function of the square of the atom fraction of the other component. This is illustrated for these same systems in Fig. 9. Some question has been raised about the

Fig. 9 Activity coefficient: liquid iron alloys.P:"?

254

Hatfield Memorial

Lectures VoL II Nx= I-NHg

-0-5~--~~~--~ '-0 0-98

0-02

0-03

0-04

005

0-96

~ 0-94

~ 0-92

~

0·00

(I-Nxt

Fig. 10

Activity coefficient of solute X in dilute liquid amalgams, at 25°C.

linearity of such plots in the region of low concentrations; extensive data in this region are available only for amalgams. As shownl" by Fig. 10, linearity persists for these systems to the lowest concentrations investigated. Generally, similar patterns are followed by substitutional solid solutions, though less data exist.

INTERSTITIAL SOLID SOLUTIONS A quite different situation is encountered in the case of interstitial solid solutions. The basic difference arises from the fact that the addition of a new substitutional atom to a crystal of necessity involves the creation of a new site; the addition of an interstitial atom does not. Put in another way, substitutional atoms may be added indefinitely, but there is a definite limit to the number of interstitial atoms that can be added on a specific type of interstitial site. The activity of an ideal interstitial solute in a binary system is found by elementary statistical mechanics to be proportional to 81 (1 - 8), where 8 is the fractional fillage of the sites under consideration: that is, the ratio of the total number of interstitial atoms in such sites to the total number of such sites. The contrast between the ideal behaviour of interstitial and substitutional solutions is shown in Fig. 11. For actual interstitial solutions, departures from ideality are to be anticipated just as in the case of substitutional solutions, already discussed. However, at sufficiently low concentration, where each interstitial atom 'sees' the same surroundings, the anticipated proportionality of activity to concentration has been amply demonstrated. These features, including the non-ideal, non-linear behaviour at higher concentration are illustrated by the data ofBan-ya et a1.20 for the iron-carbon system in Fig. 12; the same features at even higher concentration are illustrated for the Pd-Ag-H system in Fig. 13, from the data ofBrodowsky and Poeschel.F!

Metallic Chemistry in One, Two and Three Dimensions

255

9=0

Interstitial (fcc)

0

1·0

ATOM FRACTION, N2

Fig. 11

Ideal activity-compositions.

'·0

0·8

u 0

-: >~

0-6

t-

U

<

."

0·4

0·2 ".,

0

0·02

". ".

///\

Ideal,

0·04 ATOMIC

Fig. 12

"" Oc

""

e

Q

i"='e 0·08

0·06 FRACTION,

"

""

0·10

Nc

Activity of carbon in austenite, 1150°C (data ofBan-ya et al.20).

e

At sufficiently high concentration of the interstitial solute, that is, as approaches 1, we note that each vacant interstitial site 'sees' the same surroundings (full shells of matrix atoms and filled shells of interstitial sites) and hence we reasonably anticipate approach to ideal solution behaviour. However, as is apparent. from Fig. 11, the 1 term in the

e

256

Hatfield Memorial Lectures VoL II

1·0

o-e N

-.s 0

0·6

~ :x: 0

0·2

0

0·10 ATOMIC

0·20 RATIO

0·30

0·40

HIM

Fig. 13 Activity of hydrogen In Pd-25%Ag alloy, 75°C (data of Brodowsky and Poeschl-"). expression for the activity gives rise to an anticipated terminal behaviour dramatically different from the limiting Raoultian behaviour of liquid and of substitutional solutions. For this reason, experimental investigations in such a region are of special interest. Although very few have been reported, a recent investigation by Wriedt22 of the y' or 'Fe4N' phase of the iron-nitrogen system at 500°C appears of sufficient precision to check the terminal behaviour of an interstitial system. These data were obtained by weighing the Y'-phase after equilibration with various NH3-H2 atmospheres, which established the activity of nitrogen. The results of many earlier investigations of low and intermediate concentration ranges of the y field of this system are shown extrapolated to 500°C on the left of Fig. 14. The tangent (dashed) at low concentration is extrapolated from the well established solubility of nitrogen gas in y-iron; the two open circles are extrapolated from phase diagram data23 obtained by the use of NH3-H2 atmospheres. Since the iron atoms in the Y'-phase are arranged on an fcc lattice just as in the y-phase, it is appropriate to represent the activity of nitrogen (as well as other thermodynamic functions) as a continuous function of composition. Hence a smooth curve is drawn in this figure, connecting the extrapolations of the older data with Wriedt's new data for the y'-phase and reflecting the miscibility gap. The new data are shown on an expanded scale on the right of this figure. It should be noted that site fillage appears to be approached at NN = 0.2 corresponding to the composition Fe4N; this composition corresponds to nitrogen occupancy of only one quarter of all the octahedral interstitial sites: an occupancy most easily visualised as corresponding to fillage of the octahedral central sites of the fcc unit cells. The rapid rise of the activity as this composition is approached is certainly a dominant feature of this figure.

Metallic Chemistry in One, Two and Three Dimensions

257

500 40

-

0200 ~N

Fig. 14

Fugacity of nitrogen in fcc iron lattice,

soooe (Refs 22 and 23).

Let us now see whether this rapid rise conforms quantitatively to the ideal behaviour previously described. We note first that S in the region investigated is so near 1 that the ideal activity may adequately be written as proportional to 1/(1 - S) [instead ofS/(l - 8)]. Further, in such a limited composition range, (1 - 8) is adequately taken as proportional to the difference in composition between that of the actual specimen and that of Fe4N expressed in percent nitrogen. Thus a plot of 11aN versus %N is anticipated to be linear with intercept corresponding to the composition Fe4N. It is seen from Fig. 15 that such is

e

0-98

0-99

/'f)

%

a.Z

_ N (\1-

M:I: Q

"Z

S

y/_e:

Equilibrium

o~~~~~~~~~~~~~--

5 -80

Fig. 15

NITROGEN,

5-85 %

5-90 (Rz4N)

Linear reciprocal relation of activity to composition; 'Y' (Fe4N), Wriedt22).

soooe (data of

258

Hatfield Memorial Lectures VoL II

indeed the case. It is thus demonstrated that, within their range and precision, these data support the limiting law, near 8 = 1, that the activity of the interstitial, nitrogen, is proportional to Sf (1 - 8), under the assumption that only the centres of the unit cells are capable of occupancy. Recent data of Grabke-" do reveal a slight upward departure from the straight line of this figure near its lower terminus, thus indicating some occupancy of the other three quarters of the octahedral interstitial sites. Accentuation of this feature at higher nitrogen activity cannot be investigated for, as indicated in the figure, a new phase, known as E, forms if the activity is raised beyond the region of investigation. Another type of attack has been made on this problem. Schwerdtfeger et al.25 have determined the rate of growth of the y'-layer on iron exposed to various ammoniahydrogen mixtures ..Their analysis of these kinetic data does indeed indicate that a substantial fraction of the diffusive flux does occur via these 'normally empty' sites (Fig. 16). We must thus conclude that nitrogen occupancy of these 'abnormal' sites, though difficult to detect by the equilibrium measurements, is responsible for a major part of the diffusive flux.

5·80

Fig. 16

5·85

5·90

Relative fluxes of N via normal (n) and abnormal (a) sites through "/' (Fe4N), soooe (data of Schwerdtfeger et al.25).

ADSORPTION ON SURFACES AND INTERFACES Let us turn our attention now to surfaces and interfaces and the adsorbed atomic layers thereon, which we may regard as two-dimensional solutions. This subject has received much attention over many decades. However, quantitative data in the range of major metallurgical interest are relatively recent and still very meagre in many respects, in spite

Metallic Chemistry in One, Two and Three Dimensions

259

of the use of advanced techniques, such as field-ion microscopy, Auger spectroscopy and low energy electron diffraction. A major turning point in the thinking of many of us was occasioned by the findings (Fig. 17) of Holden and Kingery.F" and subsequently of Kozakevitch and Urbain.V that the surface tension of liquid iron is dramatically reduced by the presence of relatively small amounts of oxygen. OXYGE N wt -0/0 --.:O:......-·OO:....,.:...-5_0,.:·O_1 I

1800

r---__

O·--,O~O_1__

::.....,::...~~

E

~en ~1600

\J

Z

o

~

w tW

Y

Ii:

1400

1200

a: ~ 1000~--~ -8

Fig. 17

• Ha Iden and Kingery 0

Kozakevitch ~ -7 -6

and Urbain ~

Surface tension of liquid-oxygen

~ -4

__ ~

__ ~

-3

-2

alloys, 1550°C (Refs 26 and 27).

We may apply to these data the Gibbs adsorption isotherm, which states that the amount (~) of adsorption of the ith species per unit area is given by the relation T, = dcr/dJli' where cr is the surface tension and Jli the chemical potential (df..li= RT dln ai' where ai' the activity of oxygen, is here measured by the bulk oxygen concentration). It is thus found that at a bulk oxygen level of a few hundredths of a percent, well below that necessary to produce the oxide, the surface concentration of oxygen is about 1015 atoms/ cm-': a number essentially equal to the number of iron atoms/ cm-' of surface. The entire adsorption isotherm, so derived, is shown in Fig. 18. The dashed line represents the position of a' Langmuir isotherm [a DC ',9((1 -9)], the proportionality constant being chosen so 'as to fit: the data in the low concentration region. It will be noted that this relation is exactly the same as that previously cited for an ideal three-dimensionalinterstitial solution. This is hardly surprising, since Langmuir' s derivation was based on a fixed number of surface sites just as we previously assumed a fixed number of interstitial sites. Since we are now fully aware of the departure of three-dimensional substitutional or interstitial solutions from ideality, the departure of these observations for" a 'twodimensional ~olution fro~'the ideal dashed line need hardlysurpriseus, Surface chemists usually talk of such non-ideality in terms of 'lateral interaction', but it seems to me more reasonable to take the ratio of observed to ideal activity and speak of this ratio as playing the role of an activity coefficient. Unfortunately neither these data nor any of the surface adsorption data at elevated temperature appear of sufficient precision over the entire range to warrant refined treatment in the manner of three-dimensional solutions.

260

Hatfield Memorial Lectures VoL II

lJJ U

~ o-e

u, a: ::>

V)

~

w e ~

"

0·6

a: W

>

8

0·4

..J

~

Z

0

~U

0·2

,,'" --Ideal,·'. On

// /

,

/

I

/ I

/

/

/

-- -------Te e

- 0

/

"

I I

« a: lL.

0

0·01

0·02 0·03 0·04 wt -°/. (in bulk)

OXYGEN

Fig. 18

005

J

Oxygen adsorption isotherm for liquid iron at 1550°C.

As just noted, monolayer coverage is approached rapidly in the region where the partial pressure of oxygen is of the order of 10-8 atm. This strong chemisorption is clearly quite different from the physical adsorption usually measured at low temperatures and frequently explained adequately in terms of the well known Brunauer et al. 28 isotherm as illustrated-'? in Fig. 19. Here the adsorption near Po (the vapour pressure of liquid argon) may be viewed as approaching condensation. The oxygen adsorbed on liquid iron is certainly very remote from liquid 02 and must be viewed in terms of the strong chemical bonding of oxygen atoms to iron. Using the null creep method, Hondros-? has measured the surface tension of solid ironphosphorus alloys. His results, shown in Fig. 20, reveal a marked effect of phosphorus in lowering the surface tension. Applications of the Gibbs adsorption isotherm discloses a levelling off of adsorbed phosphorus at a coverage of about one atomic layer when the bulk phosphorus content exceeds 0.1%. By the same technique he has also measured adsorption at the grain boundaries of these alloys and found them to become essentially 'filled', but only by about half a monolayer, at about the same bulk concentration. Direct measurement of surface adsorption on metals at elevated temperature is confronted with the problem introduced by solubility of the adsorbed species in the matrix. Attempts to minimise this interference by the use of fine powder inevitably encounter another difficulty: the indeterminacy of surface area, compounded by the inevitable sintering. However, Cabane-Brouty-'! was able to circumvent these difficulties in her investigation of the adsorption of sulphur on silver, by use of a radiotracer technique. The low solubility of sulphur in silver and the high absorption of the ~-rays by the matrix reduced the interference by dissolved sulphur to an insignificant amount. Her isotherms for adsorption on each of three different crystallographic planes of silver are shown in Fig. 21. For the sake of clarity, points are shown only for adsorption on the 110 plane.

Metallic Chemistry in One, Two and Three Dimensions

261

o P I Po

Fig. 19

Adsorption of argon on rutile, 85 K (data of Morrison et al.29).

PHOSPHORUS

Fig. 20

J

./.

Influence of bulk phosphorus content on the surface energy of iron (from Hondros30) .

262

Hatfield Memorial Lectures Vol. II

o

'"~

)(

120 v

;n01 E

7·5

Fig. 21

Chemisorption

10'0

of sulphur on Ag {111}, {1OO}and {11 O} planes at 400°C (data of Cabane- Brouty ").

Here again, as in the preceding cases cited for adsorption on iron, we see these features of the activity-surface concentration plot: (a) an essentially linear relation at low activity; (b) an approach to a plateau or limiting surface concentration, corresponding essentially to monolayer coverage at high activity (in this case shown to be nearly the same for all three types of surfaces) and (c) anomalous or irregular behaviour at intermediate compositions. There is also strong indirect evidence of pronounced adsorption of the non-metallic elements, expecially oxygen, on metallic surfaces. Pronounced faceting of copper and nickel occasioned by slight traces of oxygen in relatively high vacuum was shown by Sundquist+' (Fig. 22); such faceting almost completely disappears, leaving spheroids when

Fig. 22 Faceting of copper and nickel crystallites occasioned by the presence of a trace of oxygen (from Sundquist=').

Metallic Chemistry in One, Two and Three Dimensions

263

these small crystals are annealed in dry hydrogen. Brenner=' has shown by field ion microscopy (Fig. 23) a similar phenomenon in the case of iridium. Many observations have been made of the striations on iron alloys when annealed in wet hydrogen. An example= is shown in Fig. 24; these striations do not appear on annealing in dry hydrogen and in fact there is a critical oxygen pressure for their appearance.v>

Fig. 23

Reversible thermal rearrangement of iridium on addition and removal of oxygen at 700°C; sequential treatments and micrographs (from Brenner=').

Fig. 24 Thermal faceting of zone refined iron after annealing in wet hydrogen (PH 0/ PH = 0.035) for 24 h at 1000°C. (a) photomicrograph using oblique light; (b) scanning electron micrograph. 34 2

2

264

Hatfield Memorial Lectures Vol. II HETEROGENEOUS REACTION RATES

It is now quite clear that the rates of a great many and probably the majority of reactions of metallurgical interest are controlled by transport phenomena, i.e. primarily by the conduction of heat and diffusion of matter. Methods of treating these frequently complex phenomena have been adopted by metallurgists from the chemical engineers. However, it is equally clear that there are metallurgical reactions conrolled primarily by chemical reaction rates at surfaces and interfaces and hence strongly dependent on adsorption thereon. One of these reactions recently investigated by Turkdogan and Martonik-v is the decarburisation of austenite by dry hydrogen. As shown in Fig. 25, when the gas flow is sufficiently rapid, the rate no longer depends upon the flow. Diffusion both on the gas phase and in the Fe-C strip was sufficiently fast that the overall reaction, C + 2H2 --7 CH4, was controlled by the chemical reaction rate at the surface. To treat this effect quantitatively let us consider first the two basic relations of the absolute reaction rate theory: (a) that the specific rate is proportional to 8:1:' the very small fractional occupancy by the activated complex, and (b) that there exists thermmodynamic equlibrium between the reactants and the activated complex. Thus, we may write specific rate

oc

(1)

8:1:

and assuming that the activated complex involves only one carbon atom and x atoms of hydrogen a:I:/acP

H2

x/2

= K:I:

(2)

Clearly, in order to use these two equations, we need a relation between 8:1: and at. We adopt the set of relations used previously expressing the activity in terms of fractional surface site occupancy (assuming that an activated complex occupies only one site) (3) where 8 = L8i and in this case 8 = bc. By combination find, specific rate

oc

8 cPH2 x/2

oc

of equations (1), (2) and (3), we

acPH2 x/2 (1 - 8)

(4)

Bearing in mind that we expect this relation to be valid only for 8c near 0 or near 1, we might thus anticipate the measured rate, at constant PH2' to be a linear function of carbon activity in the low carbon region and to be independent of carbon content in the high carbon region if there is strong adsorption of carbon there. Figure 26 shows that this is indeed the case, and we may thereby infer strong adsorption of carbon on the surface of Fe-C alloys of high carbon content. It is of some interest to note also that by perfonning similar experiments at different hydrogen pressures the exponent in equation (2) is found to be about 3/2, indicating that the activated complex includes the radical CH3.

Metallic Chemistry in One, Two and Three Dimensions

265

cttvt» .,y

o,e 0

6 V

B 16 23 30 60

z

o co

0::

«

U

I

0'5t

,, \ Diffusion \ control \~ \

\

o

300 TlME,min

Fig. 25 Effect of gas velocity on the rate of decarburisatoin of austenite (0.56 mm thick strips) in dry hydrogen at 1140°C and 0.96 atm (from Turkdogan and Martonik=").

16~--~----~--~~--~----~--~

o

01

0-2

0·3

Clc,ACTIVITY OF CARBON

Fig. 26

0·4

0-5

06

RELATIVE TO GRAPHITE

Decarburisation rate of Fe--1.S%C alloy (0.56 mm thick strips) in dry hydrogen at 1140°C and 0.96 atm (from Turkdogan and Martonik=).

The even stronger adsorption of oxygen on an iron surface may reasonably be expected to exert an even more pronounced effect on the kinetics of surface reactions, such as the transfer of nitrogen from the gas phase to iron or vice versa. Such an effect was indeed found by Turkdogan and Grieveson>? for the nitrogenation and denitrogenation of iron strips, illustrated in Fig. 27. It is quite clear that the diffusional process, which has been demonstrated to be rate controlling in the absence of oxygen, exerts negligible retarding effect here; this was further confirmed by sectional analysis revealing an essentially uniform concentration of nitrogen across the strip.

266

Hatfield Memorial Lectures VoL II

A

0·03

•

- - ---

PH20/PH2 0-332 0·214 -0 ·064 For

diffusion

PN2,atm

[%N]e

0·928 0·928 0·901

0·0241 0·0241 0·0237 0·0241

process

f!. I +J

~~

z

0·02

w

s (!)

t--

Z

0·01

o

25 REACTION

Fig. 27

50

75

TIME

I

100

h

Nitrogenation of iron strips (0.051 ern thick) at 10000e in N2 + H20 total pressure of 0.96 atm (from Turkdogan and Grieveson-").

+ H2 at

The quantitative treatment of these data is based on assumptions analogous to those just outlined for the decarburisation of austenite. It is further assumed that under the conditions of these experiments the surface sites are nearly filled with oxygen. The specific rate in the forward direction is then found to be proportional to aNI ao. Taking into account also the reverse reaction and integrating, the resulting equation for constant ao but changing nitrogen content is

log

%N -%N eg %Neq-%No

=-

ktl aol

(5)

where %Neq is the ultimate equilibrium nitrogen content, %No is the initial content, %N is the actual content at time t and I is half the thickness of the strip. Thus for any given oxygen activity this logarithmic function of nitrogen content is anticipated to be proportional to time. It is seen from Figs. 28 and 29 that such is indeed the case. The final criterion of the postulate of strong oxygen adsorption (eo == 1) and of the validity of the activity expressions as high requires that the slopes of these lines times I be inversely proportional to the activity of oxygen, that is, proportional to PH21 PH20' This linearity is beautifully verified as shown in Fig. 30. The rate obviously cannot continue to rise indefinitely as the water vapour content of the gas is decreased. At sufficiently low ratio of PH2/ PH20' the surface will no longer be nearly covered with oxygen atoms. If great care is taken to remove water from the atmosphere, the amount of adsorbed oxygen becomes insignificant and no longer affects the reaction. Grabke38 has investigated the desorption of nitrogen under such conditions.

e

Metallic Chemistry in One, Two and Three Dimensions

o

-1·5'-- __

o

Fig. 28

PH20/PH2

PN2,atm

O·0357 • 0·0347 A O' 0353

0 ·160 0 ·385 0·754

o

---'.......L.... ...I.-- __ 5 10 15 REACTION TIME, h

267

l.crn

0·051 0·051 0·051

--'

20

Nitrogenation of iron strips at various PN PN PH 0 essentially constant, 10000e (from Turkdogan and Grieveson>"). ;

2

'Zz

~~ I

/

2

2

o 10}_o.5

z z

~~ ~

Fig. 29

Nitrogenation and denitrogenation of Fe-N strips at various PN PH 0; PN essentially constant, 10000e (from Turkdogan and Grieveson-"). /

2

2

2

These very high rates reveal that the eventual plateau is many times higher than the highest rate recorded in Fig. 30. The exchange of nitrogen between gas and liquid iron follows a similar pattern-'? in that the rate is first order with respect to nitrogen and is inversely proportional to oxygen content in the usual range. The important and frequently overlooked general aspect illustrated by these examples is that, if one species is strongly adsorbed, then the activity of that species appears in the denominator of the expression for the reaction rate. If a reactant itself is strongly adsorbed

268

Hatfield Memorial Lectures VoL II

s:

E

C"

u

~

•••

X t-

Z

;$ z

tJ)

8

o

20

30

Fig. 30 Rate constant for nitrogenation and denitrogenation of iron from slopes of preceding figures as a function PN /PH 0 (from Turkdogan and Gricveson?"). 2

2

then, by cancellation, the rate may become independent in the case of decarburisation.

of the activity of that reactant, as

EFFECT OF ELASTIC STRESS ON INTERSTITIAL SOLUTES Before proceeding to the interaction of interstitial solutes with dislocations, I should like to make an apparent digression on the effect of imposed elastic stress on the activity of an interstitial solute at fixed composition, and hence on the stress effect on concentration at fixed activity. Gibbs,39 almost a century ago, considered the effect of elastic stress on the chemical potential or vapour pressure of a pure solid substance. The rather disconcerting result was that this vapour pressure depends on the direction, relative to the stress field, in which it is measured, and hence is unique only for the case of equality of stress in all directions, i.e. under the condition of hydrostatic pressure or tension. For the case of an interstitial solute, the partial pressure (strictly the chemical potential and hence the fugacity) is independent of the direction of measurement and is hence unique.t" This feature is readily demonstrated with the aid of Fig. 31. A volume element of the stressed solid,

Metallic Chemistry in One, Two and Three Dimensions 269 represented in cross-section by the shaded area, is maintained in any arbitrary stressed state by pressures (or tensions) applied by fluids, as indicated, in the principal directions of stress. We imagine the interstitial (but not the matrix elements) as soluble in the fluid and thus possessing a well defined chemical potential in each. We next imagine semi-permeable membranes, penneably only to the interstitial, at A and B and interconnect them with a tube filled with the same fluid. It is now seen that, if the chemical potential of the interstitial is different in the different compartments of fluid, there will be a flux of the interstitial; by suitable means this could be made to perform work cyclically and isothermally. Such a conversion of heat to work would be a direct violation of the second law of thermodynamics and hence we must conclude that the chemical potential and activity of the interstitial are the same in all compartments and hence may be regarded as properties of the stressed solid.

+ Py

A(fJ.M=~)

(~.fJ.~)

-----

Px

.....-fJ.~

fJ.~ Py

~

t Fig. 31

Schematic of thought experiment to show uniqueness of chemical potential of a mobile component in a body under stress.

In order to evaluate the effect of stress on the chemical potential and hence activity of an interstitial we imagine the interstitial solid solution to be carried through a reversible isothermal cycle as outlined in Fig. 32. The basic equation used is simply the statement of the second law for a reversible isothermal cycle, sometimes known as Moutier's theorum. This figure is essentially self-explanatory except that the following may be added: ~ is the external work done by the body per mole addition of interstitial; Wi = dwl dni may be called the partial molal strain of the interstitial; the superscript '0' designates reference to the same body in the unstressed state, conveniently taken as the reference state. The general relation thus obtained, for constant composition,

(6)

270

Hatfield Memorial Lectures VoL II

though deceptively simple in appearance leads to vast complications in elastic theory in many cases. However, it reduces to a simple expression for the case of an isotropic body upon which is exerted a uniaxial stress cr, small compared to Young's modulus,

In a./ a.o 1

1

crV./3RT

=-

(7)

1

where V; is the partial molal volume of the interstitial.

o

:c)-Initiol

stote

:$:

.

Work done (by body)

step 1: relax stresses

w (the elast ic energy)

step 2: transfer

o

dnj

~ .

~

Step 3: reapply stresses

-(w+ dw) 0

Step 4: retronsfer

(Wj + lij

dn,

-

lij )dnj

,

Sum of all work = 0 = (J.1rJ.1Y + Wi ~~) dnj

Fig. 32

Effect of stress on the activity of an interstitial, from reversible isothermal

cycle.

If we wish to express the concentration C, of the interstitial as a function of stress under the condition of constant activity, equation (7) is readily transformed, for small stresses, to

I C./C.o nIl

=

cr~

3RT

/(

dlnai dlnCi

)

o

(8)

At sufficiently low concentration the partial derivative is 1, by Henry's law, but in general it must be retained. Equation (8) may also be written if (Ci - CP)/ C, ~, as

e. 1

c» ==:. 1

-

0"

Vi / (

3RT

alnai)

»c, () = 0

(9)

In view of the controversy centred about this subject, Wriedt and Oriani+! have recently measured the effect of uniaxial stress on the solubility of hydrogen in a Pd-25%Ag alloy. This alloy composition, the temperature of 75°C, and the pressure of 100 torr were carefully selected to provide a high hydrogen solubility and other conditions favourable to precise experimental determination of the solubility change occasioned by the application of stress. The hollow cylindrical specimen, mounted in the stainless steel chamber and held in the jaws by which the stress was applied, is shown in Fig. 33. In operation, this whole assembly was immersed in an oil thermostat. The outlet tube was connected to volumetric gas equipment; the pressure was kept constant by a monostat. As the stress was

Metallic Chemistry in One, Two and Three Dimensions

271

Pd -.Ag specimen

Fig. 33

Arrangement of specimen, jaw assembly and enclosure for measuring the effect of elastic stress on hydrogen solubility. 41

varied, the evolution or absorption of hydrogen by the specimen was measured by a gas burette; the response was rapid, requiring only about 1 min; over the total range of stress covered the change in burette reading was almost 5 ern>, of which the correction attributable to the dimensional change of the chamber was only a small fraction. The experimental results are shown by the points in Fig. 34. The straight line drawn is that computed from equation (9), using the value of VH (1.90 cm-') as measured dilatometrically and the partial derivative as evaluated from the data of Fig. 13. The agreement could hardly be better. I hope you share with me an appreciation of this 'profounder beauty which comes from the harmonious order of the parts'. The growth of a precipitate such as cementite in ferrite frequently involves the development of a rather high stress field, not necessarily completely relieved by plastic flow. This is illustrated schematically+? in Fig. 35. In this case, using the equations previously developed and assuming that the particle is spherical and that none of the growth stress is relaxed, it is found that the growth stresses enhance the solubility of cementite by a factor of two at the eutectoid and by a factor of nearly 20 at room temperature. It is our belief that it is this stress enhancement of solubility which accounts for the major part of the former marked discrepancy between the solubility of cementite in ferrite as measured by

272

Hatfield Memorial

Lectures VoL II

internal friction, and as calculated thermodynamically from other available data shown in Fig. 36. Careful internal friction measurements by Swartz43 on both the graphite and cementite solubility in ferrite have now substantially reduced this discrepancy as shown. Our current best estimate of the solubility of stress-free cementite in ferrite is shown in Fig. 37.

NITROGEN INTERACTIONS IN IRON ALLOYS Several years ago we+" demonstrated the dramatic effect of cold work in enhancing the solubility of nitrogen in steel and giving rise to marked departure from simple solution behaviour at low concentration. This is illustrated in Fig. 38. In view of the effect of stress

't

E+15~~--~~--~~--~--~~--~~ i v +10 ~

c5

+5

~

~ o~----------------~~--------~

-30 -25 -20 -15 -10 -5 UNIAXIAL

Fig.34

0

+5 +10 +15 +20

STRESS,1031b/in2

Effect of stress on the solubility ofH2 in Pd-25%Ag alloy, PH., = 102 torr, 75°C (at zero stress the atomic ration HIM = 0.317).41 -

.:··f~{~~~::~i~~;ticle "1"·1· ,. I I

I I I 1

Solubility of

cementite

Distance from centre of Fe3C precipitate in ferrite

Fig. 35

Enhanced chemical potential and solubility of cementite in ferrite, arising from the indicated stresses associated with the precipitation (schematic).42

Metallic Chemistry in One, Two and Three Dimensions

Fig. 36

Solubility of cementite in ferrite (from Swartz+s).

_ •..--,.. ...

Fig.37

273

.",.

...•....

--

Complete Stn2SS rekixotlon

----

No stress reloxotlon

Solubility of cementite in ferrite in presence and absence of stress (from Swartz+").

on solubility, as just discussed, it now seems reasonable to attribute to those onedimensional entities, the dislocation cores, the nitrogen then associated with deep energy wells, and to the stress fields of the dislocations the nitrogen associated with the shallow wells. More recently, in a successful effort to take a closer look at enhanced nitrogen solubility in dislocated iron, Podgurski et al.45 have produced higher dislocation densities in Fe-Al-N alloys and observed the nitrogen solubility therein. The Fe-AI alloy strip was equilibrated with a series of ammonia-hydrogen gas mixtures. It is of more than passing interest to note that, in spite of the exceptionally high stability of aluminium nitride under these circumstances, nucleation of this nitride is very slow in an annealed specimen.

274

Hatfield Memorial

Lectures VoL II

0:.2

1°·03

~ Z

~0·02

o a:::

~ 0·01

o Fig. 38

01

0"1

0·2

Equilibrium of annealed and of deformed mild steel with NH3-H2 atm, 400°C.

This is illustrated in Fig. 39, where it is seen that precipitation did not start for 20 h at 550°C, but only when the temperature was raised to 575°C. In view of this phenomenon the homogeneous ferritic alloys could be equilibrated with NH3-H2 atmospheres at the lower temperatures without precipitation, leading to the rather surprising finding that under these conditions aluminium does not influence the equilibrium nitrogen content. In cold worked specimens aluminium nitride precipitates rapidly as very fine particles. The high stress associated with their growth gives rise to a dislocation network which provides nucleation sites for further precipitate particles, which in turn serve to pin the

Ea.

a.

Z w o o0: t-

Z

o

40

20 TIME,h

Fig. 39

Rate of nitrogen pickup by Fe-O.57%Al alloy in 11% NH3-89%H2 Podgurski and Knechtel+").

(from

Metallic Chemistry in One, Two and Three Dimensions

Fig.40

275

Fe-O.57%Al alloy after nitrogenation (from Podgurski and Knechtel+v).

dislocations, thus producing a very dense and stable dislocation network. This is illustrated by the electron micrograph of Fig. 40; at higher aluminium content the dislocation density is so great that the electron micrographs are nearly all black. Mter initial nitrogenation of a cold worked Fe-2%Al alloy, it was alternately reduced with hydrogen and renitrogenated to ensure reversibility and stabilisation of the structure. The equilibria subsequently. established between the alloy and atmosphere (aN always below that necessary to form iron nitride) are shown in Fig. 41. It was found that nitrogen below the point P could not be removed by hydrogen reduction. Since this amount of nitrogen is substantially above that corresponding to formation of the stoichiometric compound AlN from the aluminium present, a separate series of experiments was run. By use of the isotope N15 and the mass spectroscope it was found that the nitrogen corresponding to that between points Y and P was readily exchangeable, but that below point Y was not. Further, it was found from the measured size of the spheroidal nitride, about 100 A diameter, that the amount of exchangeable nitrogen corresponds to coverage of the interface of these particles by essentially a monatomic layer of nitrogen. Hence, it is concluded that: (a) nitrogen up to the point Y is indeed that in the stoichiometric precipitate AlN; (b) nitrogen between points Y and P is tightly bound as a monatomic layer at the interface between precipitate particles and matrix; (c) nitrogen between the levels P and X resides essentially at dislocation cores; (d) nitrogen above level X resides in interstitial sites, strongly influenced by the stress fields of the dislocations. Chou and Li,47 from the analysis of the stress field of an 'average' <011> mixed dislocation and application of equation (6), computed the iso-concentration contours around such a dislocation, similar to those in Fig. 42. The involved calculations leading to this figure incorporated the partial molal volume of nitrogen (V N = 5.9) as determined by Wriedt and Zwe1l48 and the tetragonal nature of the distortion produced by nitrogen

276

Hatfield Memorial Lectures VoL II

o

0·10 ON-

Fig. 41

020 3/2

PNH] , p~

, otm

030 -1'2

Equilibriation of Fe-2%A1 alloy with NH3-H2 (from Podgurski et al. 45).

atoms in a-iron. However, they did not take into consideration the fact that the concentration enhancement near the core is so great that the fractional fillage of such sites is not small as compared to unity, and hence the activity can no longer be regarded as proportional to concentration in this region. As discussed previously, it is more appropriate to consider the activity as proportional to Sf (1 - S) for the average occupancy of each site in the stress field. Li and Chou"? subsequently recomputed the enhancement of the nitrogen solubility attributable to dislocations on this basis (i.e. they used Fermi-Dirac instead of Boltzmann statistics). It turned out that the difficulty of the infinite stress at the centre of the dislocation core disappeared, and with it the need for an empirical energy of interaction of interstitial with dislocation core. The enhanced solubility could now be calculated with only one adjustable parameter: the dislocation density. The results of the computation, taking the dislocation density as 7 X 1011 em/ern>, is shown superimposed on the data of Podgurski et al.45 in Fig. 43. These are the same data as in Fig. 41. The origin has been moved up to point P of that figure to exclude the nitrogen as AlN and on the interface; the activity is here measured by the nitrogen content of annealed iron (equilibrated with the same atmosphere) to show the magnitude of the enhancement. Since only one adjustable parameter has been used for all three temperatures, the agreement between theory and experiment is surprisingly satisfactory. It would thus appear that we have strong support for the applicability of the general theory of the thermodynamics of elastically stressed solids to the stress fields of dislocations.

Metallic Chemistry in One, Two and Three Dimensions

277

2

2

Fig. 42

Concentration distribution around an edge dislocation; unit of distance: Jlb(l v) V/3n(1 - v)RT (-38 A for N in Fe, SOODC)(from Chou and Li47). ...!

0.

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+

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--- -0{)4 IRON,

0-05

Fig. 43 Equilibrium of nitrogen in annealed and dislocated iron lattice (curves calculated (from Li-Chou theory) for dislocation density == 7 X 1011 cm-2) (data of Podgurski et al.45).

HYDROGEN INTERACTION

IN IRON ALLOYS AND STEEL

Podgurski= has now extended his investigation to include the effect of dislocations and interfaces on the solubility of hydrogen in iron. The starting material was similar to that of point P of Fig. 41 except that the dislocation density, determined in the same way (by nitrogen equilibrium), was a little higher, about 1012 cnr='. The experiments with hydrogen were perfonned at -16° and +25°C. In this temperature range the solubility of hydrogen even at high pressure in annealed iron is far below that measurable by the volumetric technique employed. In fact, it was found desirable to use pressures up to 204

278

Hatfield Memorial Lectures Vol. II

atm (3000 Ib/In-') in measuring the solubility of hydrogen in the highly dislocated specimen. The results are shown in Fig. 44. The form of these curves, similar to those of Fig. 43, suggests that some of the hydrogen is at the particle interfaces and that the rest is dominantly associated with dislocation cores. Although the partial molal volume of hydrogen in iron has been established dilatometrically=! as VH = 2.2 cm3/gram atom, the uncertainty of the solubility of hydrogen in annealed iron precludes a detailed analysis of the data as in the case of nitrogen. The data indicate that the solution of H2 gas in this heavily dislocated iron lattice is exothermic to the extent of about 4 kcall gram atom; from this and the known enthalpy (+ 6.5 kcal/gram atom) of solution of hydrogn in annealed iron, we may infer an enthalpy decrease ofl0.5 kcal accompanying the transfer of 1 gram atom of hydrogen from a normal interstitial site to a dislocation core. PRESSURE,

100

60

1000

2000

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3000 4(X)() 5000

E50

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Fig. 44 Hydrogen solubility in ferrite as enhanced by AlN precipitation; dislocation density == 1012 cm=? (solubility in annealed iron is too small to show) (data of Podgurski 50). It would seem desirable to relate these new measurements to some of the older measurements on the hydrogen taken up by cold worked steel upon acid immersion. Over two decades ago, we reported'< data on the absorption of hydrogen by cold rolled mild steel bar (0.20%C) upon immersion in aqueous solutions of various acidity. The ultimate steady state concentration of hydrogen was determined directly. The permeability was determined from the rate of approach to this ultimate concentration. The ratio of this measured permeability to the standard permeability is the activity or fugacity. This standard permeability, taken from a recent review by Conzalez.Y' is derived by extrapolation from experiments at higher temperature in which hydrogen gas served as the source of hydrogen which diffused through an iron or steel diaphragm. Finally, the corresponding hydrogen pressure may be computed thermodynamically from the activity with the aid of the known equation of state. These data are shown as the points in Fig. 45 through which is drawn the topmost curve. On this same figure is superimposed a curve representing the estimated amount of hydrogen existing in the steel specimen as gas in the voids; these voids were introduced principally by cold work and are estimated from the density decrement of similar steels to

Metallic Chemistry in One, Two and Three Dimensions

279

PRESSURE,lb/in2

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5000

Estimated distribution of observed hydrogen content of cold rolled mild steel after acid immersion at 35°C.

occupy 0.2% of the steel volume. I have arbitrarily assumed that the amount of hydrogen on dislocations and interfaces is a quarter of that found by Podgurski in his severely dislocated specimen, on the basis that the dislocation density is here only about a quarter as much. By virtue of the judicious choices just described, the sum of all the hydrogen thus estimated on the one-, two-, and three-dimensional defects is thus brought to good agreement with the observations over this regrettably short range. In spite of the speculative nature of the breakdown of the observed hydrogen content, it is quite clear that (a) in the low hydrogen range, nearly all the hydrogen is associated with dislocations and interfaces, and (b) in the high hydrogen range, a substantial fraction of the hydrogen exists as gaseous hydrogen in voids which may provide sufficient stress for crack propagation or blister development. Before leaving the subject of hydrogen in steel, mention must be made of the relatively recent finding of several investigators including Williams and Nelson, 54 as shown in Fig. 46, that the presence of gaseous hydrogen even at sub atmospheric pressures has a marked effect in favouring crack propagation. Although the hydrogen enhancement previously illustrated is obviously involved, this crack propagation phenomenon and its ramifications are not yet fully understood.

CONCLUDING REMARKS The common theme of the many topics touched upon has been the pronounced

the one- and two-dimensional

effects of

phenomena occurring at dislocations and interfaces, on

280

Hatfield Memorial

Lectures VoL II

3~~'--~'~~'--~1---~'~1---~1~1---1~~

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the equilibrium and kinetic behaviour of the 'zero-dimensional' atoms in threedimensional metallic solutions. Although these areas of physical-chemical research on metals would be new to Hatfield, they can hardly be considered so today. Nevertheless, they still receive less attention than warranted. I hope I have conveyed some idea of the 'profounder beauty which comes from the harmonious order of the parts'. Perhaps the data and thoughts presented may serve as a minor stimulus for further research through their utility, if not their beauty. Some, if not most, of the topics discussed are a bit controversial. Hence the account given and conclusions reached contain rnore than a little of the subj ective element. This is perhaps inevitable and certainly not in violation of the spirit of Hatfield.

ACKNOWLEDGEMENT I wish to thank those members of the staff of the United States Steel Fundamental Research Laboratory who are responsible for research presented. The benefit of numerous discussions with many other colleagues, expecially E. T. Turkdogan and R. A. Oriani, is gratefully acknowledged.

REFERENCES 1. ]. H. HILDEBRAND: American Scientist, 1963,51,1-11. 2. LENO MATTEOLI: 'Diagramma di stato ferro-carbonio Centro AIM di Scuzio dei MetalIi.

e Ie curve TTT', 1 ed.; 1953, de

Metallic Chemistry in One, Two and Three Dimensions 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. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42.

s. S. BRENNER: Unpublished

281

work. S. S. BRENNER and]. T. McKINNEY: To be published. S. S. BRENNER and]. PEREPEZKO:Unpublished work. E. W. MOLLER et aL: Rev. Sci. Instr., 1968,39,83-86. S. S. BRENNER and]. T. McKINNEY: Appl. Phys. Letters, 1968, 13,29-32. S. R. GOODMAN et al.: Unpublished work. J. P. MORRIS and G. R. ZELLARS: Trans. AIME, 1956,206, 1086-90. R. SPEISERet al.: Trans. AIME, 1959,215,185-92. G. R. ZELLARSet al.: Trans. AIME, 1959,215,181-84. C. C. Hsu et al.: Izv. VUZ Chern. Met., 1961, (1), 12-20. K. SCHWERDTFEGERand H.-J. ENGELL:Arch. Eisenh., 1964, 35, 533-40. G. R. BELTON and R.J. FRUEHAN: Trans. AIME, 1969,245,113-117. G. R. BELTON and R.]. FRUEHAN:]. Phys. Chem., 1967,71,1403-1409. R.J. FRUEHAN: Trans. AIME, in press. L. S. DARKEN: Trans. AIME, 1967,239,80-89. E. T. TURKDOGAN and L. S. DARKEN: Trans. AIME, 1968,242, 1997-2005. E. T. TURKDOGAN et al.: Trans. AIME, 1969,245, 1003-1007. S. BAN-YA et al.: Trans. AIME, 1969,245,1199-1206. H. BRODOWSKYand E. POESCHEL:Z. Phys. Chem., 1965,44,143-159. H. A. WRIEDT: Trans. AIME, 1969,245,43-46. L. S. DARKEN and R~ W. GURRY: Physical Chemistry of Metals, McGraw-Hill, New York, NY, 1953, 372. H.]. GRABKE Ber. Bunsengesellschafi Phys. Chem., 1969 (in press). K. SCHWERDTFEGERet al.: Trans. AIME, 1969 (in press). F. A. HOLDEN and W. D. KINGERY:]. Phys. Chem., 1955,59,557-559. P. KOZAKEVITCHand G. URBAIN: Rev. u«, 1961,58,517-534. S. BRUNAUER et al.:]. Amer. Chem. Soc., 1938,60,309-319. J. A. MORRISON et al.: Trans. Faraday Soc., 1951,47, 1023-30; 1952,48, 840-47;J. Chern. Phys., 1951, 19,1063. E. D. HONDROS: Proc. Roy. Soc., 1965, 286A, 4179-498. F. CABANE-BROUTY:]. Chem. Phys., 1965,62,1056-1064. B. E. SUNDQUIST:Acta Met., 1964, 12,585-592. S. S. BRENNER: Suiface Sci., 1964,2,496-508. L. S. DARKEN and E. T. TURKDOGAN: Proc. Int. Con£ on Metals and Materials Science, Univ. ofPa., September, 1969. J. MOREAU and]. BENARD: Acta Met., 1962,10,247-251. scE. T. Turkdogan and L. J. MARTONIK: To be published. E. T. TURKDOGAN and P. GRIEVESON:]. Electrochem. Soc., 1967, 114,59-64. H.J. GRABKE: Ber. Bunsengesellschaft Phys. Chem., 1968,72,541-48. J. W. GIBBS: The Collected Works, Vol. 1, Longmans, Green, New York, NY, 1928. J. C. M. LI et al.: Z. phys. Chern., 1966,49,271-290. H. A. WRIEDT and R. A. ORIANI: To be published. L. S. DARKEN: Proc. 10th Anniversary of Foundation of National Research Institute for

Metals, Tokyo, 1966, 30-40.

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43. 44. 45. 46. 47.

J.

49. 50. 51. 52. 53. 54.

]. C. M. LI and Y. T. CHOU: Trans. AIME, 1969,245,1606.

C. SCHWARTZ: Trans. AIME, 1969,245, 1083-1092. A. WREIDT and L. S. DARKEN: Trans. AIME, 1965, 233, 111-121; 122-130. H. PODGURSKI et al.: Trans. AIME, 1969,245, 1603-1608.

H. H. H. Y.

H. PODGURSKI and H. E. KNECHTEL: Trans. A IlvIE, 1969,245, 1595-1602. T. CHOU and]. C. M. LI: Proc. Int. Symp. on the interactions between dislocations point defects, UKAEA, Harwell, 1968. 48. H. A. WRIEDT and L. ZWELL: Trans. AIME, 1962,224,1242-1246. H. H. PODGURSKI: To be published.

H. A. WRIEDT and H. W. WAGENBLAST: To be published. L. S. DARKEN and R. P. SMITH: Corrosion, 1949, 5,1-16. O. D. GONZALEZ: Trans. AIME, 1969,245,607-612. D. P. WILLIAMS and H. G. NELSON: Trans. AIME, 1969 (in press).

and

TWENTY-FIRST

HATFIELD

MEMORIAL

LECTURE

The Heterogeneity of Steel J. W.

Menter

At the time the lecture was given Dr]. W. Menter was at Tube Investments Limited, London. presented at The Bloomsbury Centre Hotel, London, on 2 December 1970.

The lecture was

Steel made in bulk from economic raw materials may exhibit objectionable characteristics, both in processing and in relation to end use, which are not found in small laboratory melts made from pure ingredients. These are caused by heterogeneities of chemical composition of which three of particular significance in tube manufacture are selected for detailed discussion: (a) tramp elements or residuals, (b) constitutional segregation and (c) non-metallic inclusions. The paper discusses the origins of these heterogeneities in the melting, refining and casting processes, some laboratory techniques for characterising their effect on steel properties, some methods used for quality assessment and control in production, and of non-destructive testing of the end product.

It is all too easy when working at the coal face of scientific discovery to think that no man has been there before you. In truth it is more often the case that the latter day research worker, with the advantages of increased sophistication in technique, increased spatial resolution for characterising the structure and composition of materials, enhanced computational facilities for delivering information and answers 'on line' which formerly would have required weeks and months of manual calculation, is merely fining up and lending credibility to notions and hypotheses advanced by his forebears in research. I became acutely aware of this truth in reading some of Hatfield's writings and of his contributions to the advancement of the science and technology of steel. The following quotation taken from his fascinating book 'The application of science to the steel industry'l (an extended version of the Edward De Mille Campbell Memorial Lecture given in Philadelphia in 1928), illustrates the point. Speaking of corrosion and acid-resisting steels he said: 'Today, there are alloy steels available which are truly rust resistant, and others which are so successfully resistant to the attack of certain acids as to have led to revolutionary changes in the construction of chemical plant. The author is old enough to remember the time when the metallurgical world was in a state of mind which, to say the least, was pessimistic as regards the possibility of producing such steels. The discoveries were made

.and subsequent to the essential discoveries came the explanation of the phenomena. 283

284

Hatfield Memorial Lectures VoL II

'Induced passivity had first been observed by J. Kier as long ago as 1790; in so much as he had observed that iron, after treatment with concentrated nitric acid, had lost the property of precipitating silver from solutions of silver salts, and was no longer attacked by dilute nitric acid. The great Michael Faraday, whose intellect was so profound, and whose insight was so great as to enable him to give us the key to many of nature's latent faculties to serve us, gave us in 1836 by insight the true explanation of such passivity. He reasoned that all known passivity phenomena are oxidation processes and visualised that "the surface of the iron is oxidised or the superficial particles of the metal are in such relation to the oxygen of the electrolyte as to be equivalent to oxidation". It was left to my friend Ulick R. Evans at Cambridge to establish nearly a hundred years later the existence of such a film by actually isolating it. Some eighteen months ago, the author visited Evans in his laboratory and there actually saw this 'hypothetical' film under the microscope. The polishing marks of the outer surface of the metal were definitely to be seen on the gossamer-like patches of the films which were floating under the objective. "Are they actually oxide of iron?" the author asked, and Evans at once caused the reaction and the films turned blue. It may be imagined how deeply impressive was this experience. Here was the complete demonstration of the accuracy of the view held by both of us, that passivity was to be explained by the existence of this film.' I have quoted at length for several reasons. Firstly, because the passage illustrates so well how we all stand on the shoulders of those who have been there before us and how so much of our work today is indeed a refinement and proving out by experiment of old ideas and empirical practices. Secondly, for the insight it gives, through the delightful account of the episode in Evans' laboratory, into Hatfield's approach to scientific discovery. Research directors and managers, harassed by requirements for cost-benefit analysis and the like, should never forget that the excitement and joy of discovery, with the personal satisfaction it brings, are still the most potent driving forces in scientific advance. Thirdly, the quotation may encourage some authors, of whom there are today unfortunately all too many, to embellish their papers by exploiting the full descriptive range of the English language. I must admit this is a fond hope, 'gossamer' would never be selected as a key word by the information scientist using computerised retrieval, and if it were, it would be put together with 'floating' to classify the paper under the heading of some obscure species of water spider, thus ensuring that it would never subsequently come to the attention of a metallurgist. Fourthly, the quotation gives me a due sense of awe in accepting the honour of being invited to deliver this Hatfield Memorial Lecture, in that my first encounter with steel was in 1950 when I used the electron microscope in conjunction with electron diffraction to establish, in collaboration with J. E. O. Mayne and M. J. Pryor.? working in that same laboratory ofUlick Evans, that the protective film on mild steel was indeed gamma iron oxide. I had not appreciated at the time that Michael Faraday had been there before me! My humility in the face of my scientific forebears and Hatfield in particular is no less deep in relation to the topic I have chosen for this address. The heterogeneity of steel

The Heterogeneity of Steel 285 ingots was the subject of a classic series of cooperative researches inspired and led by Hatfield in the last twenty years of his life, researches jointly sponsored by The Iron and Steel Institute and the Iron and Steel Federation, which foreshadowed the subsequent great extension of cooperative research within the industry in the form of the British Iron and Steel Research Association under the leadership of Sir Charles Goodeve. At the time of Hatfield's death no less than seventy-eight ingots, ranging in weight from 13 cwt to 172 tons had been sectioned and examined, and the results of this examination are on record in a series of weighty reports from The Iron and Steel Institute Heterogeneity in Steel Ingots Committee. It is not my intention to consider in detail the present state of knowledge of ingot structure, but to use the privilege which is customarily allowed to Hatfield lecturers to discuss selected aspects of my subject 'The heterogeneity of steel', including ingot structure along with others. My selection is that of a physicist interested in crystal microstructures, who twelve years or so ago was astonished to learn that there was more to steel microstructure than ferrite, pearlite, martensite, bainite, austenite and the rest, comprising alloys with tailor made distributions of carbides with improbable compositions, giving tensile strength, creep resistance, fracture toughness, etc., to specification at whatever service temperature was required. All of that type of heterogeneity I take for granted in the present context. I wish to talk about those heterogeneities which are a nuisance and hindrance to the achievement of those ideal properties which a textbook microstructure should have. I cannot treat all these nuisances, but will confine myself chiefly to three which have been my concern in directing research on tube steel in the last ten years. Here I must make proper acknowledgment to the staff of TI laboratories who have carried out much of the work I shall be describing and to friends in universities and the steel industry who have made material available for my lecture. The three topics I have chosen are the so called tramp elements or residuals in steel, constitutional segregation and non-metallic inclusions. I shall discuss the origins of these heterogeneities in the melting, refining and casting processes, some laboratory techniques used for characterising their effect on steel properties, some methods used for quality assessment and control in production, and of non-destructive testing of the end product.

RESIDUAL ELEMENTS AND THEIR EFFECT ON MECHANICAL PROPERTIES Surface Hot Shortness Although bought scrap accumulated by scrap merchants is a highly attractive source of iron for charging into steel furnaces, it is relatively uneconomic to separate out that component containing elements other than iron known to be detrimental to the properties of steel. In particular, copper and tin when present above a certain level of concentration give rise to a phenomenon known as surface hot shortness during the process of hot rolling. Transverse fissures may appear on the surface of the rolled product leading in the

extreme to a complete break-up of the surface (see Fig. 1). Since economies demand that

286

Hatfield Memorial Lectures VoL II

Fig. 1

Plan (top) and section (bottom) of mild steel billet exhibiting surface hot shortness= (x 2).

contaminated scrap must be used, attempts must be made to understand the effect in order to set tolerable limits on the concentration of offending elements in the steel. Empirical observation over many years had shown that in mild steel with percentage compositions of copper and tin conforming with a formula of the kind Cu + 8Sn < 0.4, difficulty was not encountered in practice. There is an abundance of such formulae to guide everyday practice in a steelworks, and they represent a challenge for the scientist to seek an explanation of the underlying phenomena, in order thereby to improve practice by a better appreciation of the interaction of the relevant material and process variables. Dr Melford," working in our Hinxton Hall laboratories, took up this challenge some twelve years ago by using (and subsequently playing a significant part in developing further) the new technique of scanning electron probe microanalysis. This technique, invented by Castaing and developed into a scanning form by Cosslett and Duncumb.> represents perhaps the most important development in metallography since Sorby first showed that phases in alloys could be distinguished by using the optical microscope to examine suitably polished and etched specimens. A measure of the explosive growth of its use is given by the installation of several hundred instruments throughout the world since they became readily available commercially scarcely more than a decade ago. A casual perusal of the

papers of any metallurgical journal concerned with microstructure is enough to demon-

The Heterogeneity

of Steel

287

strate how much the technique has now become an accepted and essential tool of this field of study. Returning to the problem of surface hot shortness, Melford examined normal sections taken from low carbon (0.08-0.15%) steels and first confirmed that the problem arose from surface oxidation in the heating furnace used to raise the billet temperature to about 1130°C ready for hot rolling. As transformation of the surface of the billet to iron oxide proceeded, there developed a local enrichment at the oxide scale/steel interface of the elements less susceptible of oxidation, notably copper, tin, nickel, arsenic and antimony (Fig. 2). The local concentrations at this interface region were shown by electron probe microanalyses to be as high as 7% copper and 1% tin compared with a volume concentration of these elements through the main bulk of the steel, as measured by conventional chemical analyses, of 0.20 and 0.06% respectively. The quantitative measurement by electron probe analysis of the local enrichments of the offending elements, together with other quantitative studies of synthetic alloys of Fe-Cu-Sn, Fe-Cu-Sb, etc. helped Melford to construct phase diagrams establishing the effect of tin, antimony, etc. on the solubility of copper in austenite (see Fig. 3).

Fig. 2 Scanning electron probe micrographs of normal section through oxidised mild steel surface showing subscale enrichment of residual elements= (x 245); (a) electron; (b) tin; (c) antimony; (d) copper; (e) nickel; (f) arsenic. From these he deduced enrichment factors which, if exceeded in the subscale region, would lead to the creation of a copper rich phase which was liquid at the hot working temperature causing a loss of mechanical strength and consequent surface break up under

the tensile stress of rolling.

288

Hatfield Memorial Lectures Vol. II

<,:;------------

~x

--Arsenic

10

5 NICKEL ,TIN,ANTIMONY,

Fig. 3

OR ARSENIC,wt-%

Effect of small additions of nickel, tin, antimony or arsenic on the solubility of copper in austenite at 1200°C.3b

The nature and extent of the enrichment process was found to be a function of time and temperature within the furnace as well as oxygen potential, and it became possible to treat quanitatively for the first time all of these variables, together with those of composition of the steel and thus to obtain more flexibility than had hitherto been possible in adjusting process parameters to avoid trouble. As one might expect, this flexibility is limited and while it is sufficient to enable us to live with tramp elements today by setting upper limits to their concentration in steel melts, the long term trend in levels is in an upward direction which is enhanced by the process of internal circulation of scrap within the UK between steelworks and user. Proposals have been made in the past for slag reactions in steelmaking which will reduce the concentration of tramp elements in liquid steel, but none has yet emerged which is both effective and economic. More research should and no doubt will be done in this direction, but prudence demands a search in a totally different direction for a solution. Especially in relation to electric arc steelmaking the best hope seems to lie in the use of prereduced iron pellets as feedstock to replace at least part of the scrap charge in order to maintain tramp element concentrations at an acceptable level. The increasing availability of rich iron ores and cheap sources of fuels for reduction together with the economies in sea transport costs from distant places made possible by the introduction of bulk carriers, lend support to the belief that suitable iron with negligible tramp element content may in due course be available at an economic cost compared with scrap. Although the price ton for ton of iron may be greater, one may anticipate reductions in conversion cost stemming from the use of a physically more

The Heterogeneity of Steel

289

homogeneous raw material used in a furnace specially designed to take advantage of this, so that the resultant ingot cost may not be markedly different.

Grain Boundary Segregation of Residual Elements So far, my discussion of the effect of residual elements on steel properties has been confined to the problem of surface hot shortness. It is also possible for some of these elements to reduce the bulk ductility of steel with serious consequences both in hot working and in subsequent use. Among the most striking effects are those arising from very small bulk concentrations of certain elements such as antimony and phoshorus which segregate to grain boundaries to such a degree as seriously to reduce intergranular cohesion leading to intergranular fracture under stress at comparatively small extensions. No phase separation can be seen associated with this phenomenon, even using the elctron microscope at the limit of its resolving capabilities, and it is surmised that the segregation may be no more than a few atomic distances in extent normal to the grain boundary surface. Here again, as with surface hot shortness, we have scientific knowledge based on inductive reasoning and works practices based on accumulated experience. Advances in observational and measurement technique may well lead to a deeper level of understanding with concomitant advantages in improved practices. Several developments currently emerging from physical research deserve consideration.

PHYSICAL TECHNIQUES FOR STUDYING GRAIN BOUNDARY SEGREGATION Electron Probe Microanalysis (EMMA) The electron probe analysis technique to which I have referred earlier is limited in its spatial resolution of element concentrations by the unavoidable scattering of the electrons responsible for X-ray excitation into a volume of material with linear dimensions determined by the energy of the incident beam comparable with the probe diameter. The limitation persists with a solid specimen even if the probe diameter is significantly reduced, but the advantage in resolution obtained with a smaller probe can be realised" if, instead of a solid specimen, the material to be examined is prepared in the form of a very thin film (see Fig. 4), similar to that commonly used in the so called thin film method developed by Hirsch and others for studying the microstructure of materials in the transmission electron microscope. Indeed, it has proved possible to modify such a microscope by including an additional lens (a mini lens) in the illuminating system to permit the formation of an electron probe 0.1 Jlm or less in diameter, and to add crystal spectrometers for measuring characteristic X-ray emission from the specimen, in a manner now making it possible to obtain quickly, and with no doubt about the location of the point of

interest, a transmission image, quantitative chemical composition

and crystal lattice

290

Hatfield Memorial

Lectures VoL II

structure (by electron diffraction) from a preselected area of a specimen less than 0.1 J..Lm in extent. This new facility (EMMA 4) (see Fig. 5) for microstructural studies has been incorporated in the AEI EM 802 microscope through the work of Cooke and Duncumb.? The potential of the instrument for studying compositional heterogeneities on a very fine scale is shown by the solute depletion of zinc in the vicinity of a grain boundary in an Al-26 wt-%Zn alloy after heat treatment ofl h at SOO°C, followed by 4 h at 240°C (jacobs") (Fig. 6).

Field Ion Microscope Atom Probe An alternative approach to quantitative studies of segregation effects at the atomic level which may prove even more fruitful, if somewhat narrower in scope, is the field ion microscope (FIM) atom probe first developed by Muller and Panitz,? and later in a form particularly suited to metallurgical needs by Brenner and Mclxinney!" and by Turner and Southon.U You will recall that the simple field ion microscope originally due to Muller permits the direct imaging of the atomic structure of the specially prepared tip of a metal

Fig.4

Computer

plots of scattering of20 kV electrons in aluminium targets of thicknesses indicated; sideways scattering is least in thinnest foil. 6

The Heterogeneity of Steel

Fig. 5

EMMA 4-AEI transmission electron microscope facility.

291

with electron probe microanalysis

or alloy specimen by causing inert gas ions (typically helium or neon) to move from adsorption sites on the surface, under the influence of an applied electric field, to fall on a fluorescent screen where their impact creates a highly magnified image closely reflecting the spatial distribution of metal atoms on the surface of the tip. In the atom probe modification of this technique, the applied field is augmented until it causes the surface atoms themselves to evaporate and move as charged ions towards the screen which now has a small hole in it through which ions may pass into a time of flight mass spectrometer. By manipulating the metal tip, any preselected area of the surface image may be caused to lie over the hole in the screen (Fig. 7), which may be so small that a particular atom from a particular point on the surface may be caused to pass through the hole so that its mass and hence its atomic identity may be established by means of the mass spectrometer. One can scarcely imagine a more direct method of observing and measuring heterogeneity at the atomic level, and we look forward with anticipation to its use on problems of grain boundary segregation.

292

Hatfield Memorial Lectures VoL II I

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I

(b)

Fig. 6 Study with EMMA 4 of zinc depletion adjacent to grain boundary in AI-26 wt-%Zn alloy" (a) transmission electron micrograph, 1 h, SOO°C;+ 4 h, 24DC; (b) plot of Znl AI·count rate ratio measured at numbered points on micrograph.

Fig. 7

Field ion emission image of tungsten point. Circular aperture centre indicates area selected for atom analysiswith mass spectrometer.l '

The Heterogeneity of Steel

293

Auger Spectroscopy A third method offering the possibility of establishing the presence and perhaps the local concentration of a segregated element on an exposed fracture surface is that of Auger spectroscopy (sometimes also called electron spectroscopy). The possibility was realised some 17 years ago by Lander-? but it is only within the last three years, following the work of Harrisl-' and others, that interest in the technique has revived. The basis of the technique is to measure the energy spectrum of electrons ejected from a surface under the impact of a monochromatic incident electron beam. The energy spectrum of these so called Auger electrons is uniquely characteristic of the surface atoms present. Because of the relatively low energy of the excitation electrons (up to 1 kV) the technique is particularly suited to the study of elements located within the outermost atomic layers of the surface. Marcus and Palmberg-" have studied an AISI 3340 steel (2.2%C, 1.88%Cr, 3.44%Ni, 0.03%Sb) using this method and have established unequivocally the segregation of antimony at grain boundaries during an embrittlement heat treatment and that the segregation is limited to the region close to the grain boundaries. Figure 8 shows the comparative Auger spectra of embrittled and unembrittled samples.

Antimony

e

Fe

(a)

Fe

Unembrittled surface

o

(bJ

600 VOLTAGE,

Fig. 8

Fe Fe

eV

Auger spectra of fracture surface of AISI 3340 steel.!" (a) embrittled; (b) unembrittled.

294

Hatfield Memorial Lectures VoL II HETEROGENEITY OF CAST STEEL

Of all the stages in steel production from raw material to finished product, the least understood is undoubtedly the transition from liquid to solid, and yet it is perhaps the most important for it is in this process that are created those heterogeneities which by their persistence, albeit in modified forms, throughout the heating and mechanical working which follow solidification, may cause serious difficulties in processing and property deficiencies in the end product. Much may be learned about the origin and nature of these heterogeneities by metallographic studies of sectioned ingots and continuously cast blooms. I have referred already to the extensive studies of this type carried out under the auspices of the Heterogeneity Committee. Interest in this approach seemed to wane following Hatfield's death until the arrival of the electron probe analyser which provided the means of mapping the chemical heterogeneity of the ingot in much greater detail than was formerly possible. Before discussing some of the results of this later work, which has contributed materially to improving steelmaking practices, I would like to emphasise that in relation to many important features of steelmaking, the study of the sectioned ingot has all the limitations of a post mortem investigation. We badly need new techniques for obtaining nlore information on the physical and chemical dynamics of the solidification process. Research workers in universities and elsewhere could bring great benefit to the industry if they would give more attention to this most fundamental aspect of process metallurgy. The heterogeneity found in solid steel may be broadly classified into three types: (i) void (ii) local differences in composition (iii) non-metallic inclusions.

of a constitutional nature

Any combination of these may occur together at some point in the ingot, but it is convenient to treat them separately in the present context.

Void Voids are generally found to be innocuous if they are totally enclosed with no access to the external atmosphere, and provided additionally that they are not associated with nonmetallic inclusions, they disappear in the course of mechanical working. It is interesting to note in passing that the so called 'central looseness' or microvoid commonly found in the central region of continuously cast steel has no adverse effect on the bore surface quality of rotary pierced tubes made from concast billets, provided the central region of the billet is free from non-metallic inclusions.

Constitutional Segregation Four distinct processes are at work in the solidification of an ingot:

The Heterogeneity of Steel

295

(i) (ii)

formation of a thin chill layer immediately adjacent to the ingot mould wall columnar or dendritic growth upon the chill layer of crystals, elongated roughly normal to the ingot mould wall (iii) accretion of solid on the bottom of the mould as a result of the accumulation of crystals falling through the central pool of liquid metal (iv) a final filling in of the central region resulting from the meeting together of growth fronts from sides and bottom influenced by the liquid circulation process in the ever diminishing central pool.

It is not surprising that the continuously changing physical conditions of these processes, compounded with the chemical complexities of phase relationship in alloys with a finite freezing range, lead to variations in chemical composition throughout the frozen mass. In some alloys, e.g. low carbon rimming steel, the resulting heterogeneity may be turned to advantage but it is more often the case that it gives rise to highly undesirable effects. A notable example of this is the 1%C, 1.S%Cr alloy used for bearings, of which TI makes many thousands of tons in the form of tube supplied to bearing manufacturers. This steel has, therefore, been the subject of continuing study over many years in our laboratories (Melford and Doherty,"> Melford and Granger!"). Several varieties of segregation may be distinguished, assocciated with the different stages of solidification outlined above. There is first a process of microsegregation (see Fig. 9) occurring within the last-tofreeze pools of molten alloy entrapped within the interstices of the dendritic growth zone. This can be sufficiently severe to produce a second phase in the interdendritic regions in the form of a eutectic chromium-rich carbide with a melting temperature of l130°C ± 10°C. Since the preferred hot working temperature of this steel is in the same region, it will be seen immediately that trouble may be expected unless this phase is

Fig. 9

Central portion of a 5 ton ingot of 1%C-l.5%Cr steel showing carbide segregation in interdendritic pOOIS.17

296

Hatfield Memorial Lectures VoL II

eliminated. The diffusion distances required to restore the local uniformity of composition are fortunately small with this type of microsegregation, so that it can be readily eliminated by a soaking treatment of 4 h at 1150°C before rolling. More difficult to deal with, however, are the phenomena of gross segregation in the form of the well known A and V segregates associated with the later stages of solidification. Their influence on tube manufacture is shown in the following series of illustrations. Figure 10 shows the principles of rotary piercing used to form a hollow from a solid billet by cross rolling over a plug supported on a bar. Figure 11 shows a longitudinal section through a billet in the process of piercing, stopped in the piercer with the plug in position. Figure lla shows no break up of the material immediately ahead of the plug in a billet without segregation compared with Fig. lib where, due to the poor hot working characteristics of the axial segregated region, distinguished by the dark etching effect, internal fracture has occurred ahead of the plug. The latter condition leads to an unacceptable bore surface as shown in Fig. 12. This axial segregated region in the billet can be related back to the V segregates of the ingot. The A segregate also gives rise to a tube defect through an analogous effect of poor hot workability known as midwall shear and illustrated in Fig. 13. In this case, under particularly bad conditions, it is even possible to cause almost complete midwall separation, giving two coaxial 'tubes'! The unique association of chromium segregation with the region of poor hot ductility is shown by Fig. 14 which is a line scan of chromium concentration made with the electron probe analyser along a line normal to the tube bore surface. There is a sharp rise at the boundary between the unsegregated and segregated regions. Hot notch tensile tests (see below) on specimens from segregated and unsegregated billets confirm the poor hot

Roll

Fig. 10

Schematic diagram of tube hollow production by rotary piercing of solid billet by cross rolling over plug.

The Heterogeneity of Steel

297

Fig. 11 Interrupted piercing of En31 billet.J'' (a) no segregation, no fracture ahead of plug; (b) axial segregation shown by dark etching effect on extreme right, fracture ahead of plug.

Fig. 12

Longitudinally

sectioned tube obtained from En31 with axial segregation.

18

ductility associated with the former (see Fig. 15). The metallurgical situation in the segregated regions may well be more complex than a simple enrichment of chromium and carbon since other elements, for example, sulphur and phosphorus, are also found to be concentrated in these regions and the poor ductility may be worsened by burning effects either in ingot and billet heat treatment or by local temperature rises within the billet in the piercing process itself Be that as it may, there is no doubt that once this level of segregation has been frozen into the ingot there seems little practical possibility of dispersing it by heat treatment, and the lesson to be learned from these studies is that ingot

298

Hatfield Memorial Lectures VoL II

Fig. 13

Transverse

section

of tube obtained from segregation. 19

En31

billet with

midradial

casting practice must be so controlled as to avoid the segregated regions forming in the first place. Restriction of ingot size is certainly a help, although it would scarcely be economic with this particular steel to resort to the technique of 'ingot production' in the form of spheres 0.5 mm in diameter, recently announced by Stora Kopperberg.s! which neatly eliminates the severe segregation effects encountered in the traditional methods of casting highly alloyed tool steels! Teeming procedures are doubtless a factor strongly influencing solidification processes. We have much to learn about what these influences are.

o

TUBE SURFACE

50

100

150J,.lm -..

TUBE

BORE

Fig. 14 Line scan with electron probe microanalyser across boundary of segregated region in En31 steel tube showing rise in chromium concentration associated with segregation.lf

The Heterogeneity of Steel

Hot notch tensile

299

tests

100

tf!

-Good casts (free from gross segregation

80

«

w

0::
60

6

z

0

)::.

40

U :::>

0

w

a::

20

0 1040

Fig. 15

1160 TEMPERATURE,oC

1280

Comparative hot ductility of segregated and unsegregated casts of En31 steel shown by reduction in area in hot notched tensile test.20

Non-rnetallic Inclusions I turn now to consider my third type of heterogeneity, non-metallic inclusions, their mode of occurrence, their effect on mechanical properties and methods of assessing them. Figure 16 shows a typical collection of inclusions obtained by Bergh22 from a piece of silicon killed mild steel, by dissolving away all the matrix. Beautiful as they are in this state of splendid isolation, when they are back in their natural habitat, the steel ingot, they are undoubtedly the most tiresome and all pervading of heterogeneities.

Fig. 16

Non-metallic

inclusions extracted from cast silicon killed mild steel.F

300

Hatfield Memorial Lectures Vol. II

No steel is totally free from them. They occur at densities between 1012 and 1015 per ton, they are observed, characterised, counted and sized daily in their thousands and delivered annually from steelworks in the UK to the tune of several thousand tons. Inclusions are one of those rare commodities for which the customer pays more for less in the sense that there is almost always a cost penalty associated with the steelmaking methods required to reduce their incidence. Thus in controlling the occurrence of inclusions in steel, there is a balance to be struck between the cost of making steel and the extent to which the engineering properties of the product are diminished by their presence in relation both to the needs of secondary processing (rolling, forming, etc.) and end use. For example, with a product such as a stainless steel fuel element can with a wall thickness of 0.6 mm, which must reliably contain radioactive fuel to ensure no leakage into the cooling gas flowing past it, it is clearly prudent to use vacuum induction melting followed by vacuum arc remelting for preparing the starting material for tube manufacture, to reduce to a minimum the probability that there may be an inclusion extending through a significant proportion of the wall thickness, leading to the possibility of rupture under the working stresses of the reactor. By contrast, the additional costs associated with this method of obtaining extremely clean steel would be unsupportable, and in any case unnecessary, in relation to a product such as a stainless steel sink unit. Even with the latter product, however, special care must be taken to ensure reasonable freedom from gross inclusions since the severe forming operations required to obtain the finished shape may cause splitting at inclusion stringers, or lead to an unsightly appearance of the surface. The problem of living with inclusions may be broken down into several parts: • what is their origin and how may they be controlled at source? • how do they affect engineering properties? • how may their incidence be assessed for purposes of quality control of the product? I shall now consider these questions in turn.

Origin of inclusions The three significant sources of inclusions are the refractories used to line the vessels and runners necessary for melting, refining and casting steel, the slag used for refining and the reaction products arising from the additions required to complete the deoxidation of the liquid steel. It used to be thought that the inclusions observed in the solidified steel could each be uniquely associated with one or other of these sources. Charles and SahnonCox23· undertook an extensive survey of a complete midsection slice from a 3 Y2ton killed 0.2% carbon steel ingot using the electron probe microanalyser and found that although most of the non-metallic matter in the ingot was endogenous (i.e. resulted from deoxidation reactions), many inclusions contained elements which could only have arisen from furnace slag and erosion of ladle lining, e.g. the high level of calcium found in the inclusion shown in Fig. 17. The elements characteristic of these sources were generally diluted in the inclusions compared with their concentrations at source, indicating a

The Heterogeneity of Steel 301

Fig. 17

Scanning electron probe pictures of non-metallic inclusion in 0.2% carbon steel ingot, photographed with radiations indicated.F'

dynamic interaction between non-metallic material of differing origins in the course of removing steel from the furnace to the ingot mould. Pickering=' and his colleagues have recently reported a comprehensive investigation into the nature of the changing population of non-metallic inclusions present in a steel melt throughout the various stages of manufacture which re-emphasises the relative importance of exogenous material as a source of objectionable inclusions. The relative volumes of typical inclusions from a variety of sources found in ingots of high carbon and low alloy steels made to a double slag practice in a basic electric arc furnace are given by Pickering-" as follows:

Type of inclusion Alumina, spinel and CaO.6Al203 (other than clusters) Other calcium aluminates Secondary deoxidation products (Si killed steel) Primary deoxidation products (Si killed steel) Erosion silicates (Al killed steel) Erosion silicates (Si killed steel)

Diameter, J..lm 5 27 32 49 64 107

Relative volume 1 160 260 940 2100 9800

302

Hatfield Memorial Lectures VoL II

Insofar as the larger inclusions may be expected to cause more trouble in mechanical working and impair product properties it is clear that the last three of the sources listed are most important. Although the incidence of the primary deoxidation products may be minimised by a carefully controlled deoxidation practice, a very clean bath of steel can still be seriously spoiled by pick-up of large inclusions by erosion of launder, ladle or runner refractories in the process of teeming, or by the action of circulating currents within the ingot, drawing down scum or parts of hot topping tiles which then become entrapped in the solidifying steel. There is some substance in the furnaceman's old complaint that a bath of first class steel can be ruined in the casting bay! These studies which I have mentioned, along with many others, particularly those of Kiessling and his colleagues, all using the electron probe analyser as an essential tool, have established the equivalent of the naturalist's 'fauna and flora' in the form of an atlas of non-metallic inclusions prepared as a series of special reports for The Iron and Steel Institute by Kiessling and Lange.25 With the more precise indications of the origin of the many species of inclusions provided by this systematic work it is now becoming possible to attack the problem of eliminating inclusions at source in a more systematic way, and to establish the areas of further research which might be particularly fruitful, e.g. to improve understanding of the mechanisms of refractory erosion with a view to enhancing their resistance to attack by liquid steel. In addition we are rapidly gaining a better understanding of what are the real events, as distinct from supposed events, occurring in works practice when vacuum degassing, pouring through slag, holding ladles, etc. so as to determine and control works practice on a more scientific basis than was possible before. Thus one may expect in the future a continuous improvement in the cleanness of common steels without concomitant cost increases, but for the high levels of cleanness required for exacting end uses it will still be necessary to incur the additional costs of secondary refining processes such as electroslag remelting or vacuum arc remelting.

Some effects of inclusions on mechanical working properties While it is useful and important to obtain a comprehensive understanding of the origin and nature of all the non-metallic inclusions in a steel ingot, it is essential in practice to adopt a highly selective approach to the mitigation of their influence on steel properties. In mechanical working this selectivity is related to the particular mode of deformation to which the steel is subjected and in this respect the manufacture of tubes by rotary piercing of solid billets serves as a good illustration. It is apparent from the illustrations of this process in Fig. 10 that the bore surface of the tube is created from the material in the central region of the billet. Moreover, the very severe deformation involved in creating the bore by the action of cross rolling before and over the plug makes it especially likely that inhomogeneities in this region will initiate internal fracture, resulting in fissures and laps on the bore surface. Thus in this particular instance it is the inclusion population in the vicinity of the billet axis which critically governs hot ductility. Cottingham.v= therefore, developed a hot notched tensile test in which a test specimen was cut from the tube

The Heterogeneity of Steel

303

billet in a shape (Fig. 18) which permitted a searching test of the hot ductility of a substantial volume of the central region of the billet, and found that the reduction in area observed in the test correlated extremely well both with the incidence of bore defects in tubes made from billets from which the testpieces were taken and with the volume fraction of inclusion material in the central region of the billet measured optically on a polished transverse section. The fibrous appearance of the fracture surface (Fig. 19) showing poor ductility resulted directly from inclusions several em in length which had been elongated into needles from a globular form in rolling from ingot to billet. The appearance of such inclusions in the original billet is shown strikingly by a deeply etched transverse section in Fig. 20. The manner in which inclusions deform in the process of rolling in itself can be a controlling influence on whether or not they cause undesirable defects. Our investigations of continuously cast billets as raw material for tubemaking (see, for example Holden et al.27) illustrate this point and add further weight to my earlier remarks about the need for selective consideration of the effect of a given inclusion population. Tube billets (0.08-0.13%C) were continuously cast as 150 mm rounds with the aim of using these in place of rounds of the same size rolled from ingots, thereby gaining economic benefit both from the yield improvement from liquid to solid steel characteristic of continuous casting and from the elimination of the primary and secondary rolling costs associated with the reduction of ingots to billets. The hot working of the tubemaking processes was

Fig. 18

Notched tensile specimen cut from 4 in billet.26

304

Hatfield Memorial Lectures VoL II

Half of hot. notched tensile test specimen after fracture at 1100oe; the fibrous appearance of the fracture surface is caused by long needle-like non-metallic inclusions in central region of billet. 26

Fig. 19

Fig. 20 Deep etched transverse section of mild steel billet showing needle-like nonmetallic inclusions in central region; some globular inclusions in cast steel can become

highly elongated during rolling. 26

The Heterogeneity of Steel

305

adequate to break down the cast structure of these billets to give an entirely acceptable microstructure but an unacceptably high incidence of bore defects was encountered (Fig. 21), particularly after a fine finishing operation, honing, required in the production of tubes for hydraulic applications. The overall cleanness of the steel was very good at less than 40 ppm oxygen with a volume fraction of inclusions in the central region of the bar of only 0.01 %. Metallographic studies showed that even this low level could indeed cause bore defects and furthermore, that the offending inclusion material was smeared out over a rather small volume into a lenticular shape (Fig. 22), presumably because of the axial and circumferential shearing strains arising in the rotary piercing process. It appeared that although the original inclusion or inclusion aggregate had been fragmented by the severe stresses and strains of hot working, the fragments were still close enough to cause generalised local fracture of the steel. It was concluded that primary elongation of a concast bloom before rotary piercing might improve matters. The inclusion material would be reduced in transverse section and dispersed axially not only because of the greater elongation but also because of the absence of circumferential shear in conventional rolling. So indeed it turned out with steel made to a similar practice to that used before but in the

Fig. 21

Defect in surface of honed bore tube caused by presence of non-metallic inclusion near central region of billet. 27

Fig. 22

Successive section at depths indicated through bore defect shown in micrograph on extreme left; section parallel to bore surface28 (x 8).

306

Hatfield Memorial Lectures V01. II

form of260 mm octagons rolled to 150 nun rounds before rotary piercing. No significant defects have been found in honed bore tubes using billets made in this way, notwithstanding that the inclusion content of the central region of the 260 rnrn octagon was in fact several times greater than that of the 150 mm rounds. A further significant conclusion from these studies was that the frequency of offensive inclusions along the centre line of the concast billet was such that there was a relatively high probability of finding at least one in a length of say, one metre of finished tube, which was about the length required by the customer. Similar bore defects are sometimes encountered in tubes made from ingot steel, also attributable to non-metallic inclusions, generally much larger than those found in concast steel. However, the frequency of occurrence near the critical bore region is very much smaller so that the proportion of metre lengths to be rejected is also very much smaller. Thus the more uniform dispersion of non-metallic material in concast steel compared with ingot steel gives no advantage in this particular application, although in others it does. These illustrations taken from some of our own experience over the very limited field of the hot ductility of carbon tube steel demonstrate the immense complexity of the problem of the assessment of the relevant inclusion content of steel. Size, degree of dispersion, location, deformation characteristics of the inclusions together with the deformation modes of the steel matrix, all of these must be considered in trying to decide what can and cannot be tolerated. It is not surprising that there is as yet little agreement among metallurgists and those responsible for quality assurance about the best methods of assessment. Undoubtedly we require more detailed studies of the mechanics of individual events involving inclusions and their deformation. The recent literature shows evidence that some such studies are in progress. For example, Charles and Maunder-? and Charles and Uchiyama 30 have made useful metallographic studies of the deformation characteristics of silicates and sulphides in 0.2%C steel under hot working conditions. Kiessling and N ordberg-'! have started to apply the concepts of fracture mechanics to crack initiation by inclusions in an effort to distinguish between offensive and inoffensive inclusions. Brooksbank and Andrews+' have made quantitative studies of the 'tessellated' stresses arising from differences in thermal contraction between inclusion and matrix on cooling from hot working temperatures and their possible influence on the fatigue life of bearing steels through their tendency to cause local matrix yielding and crack initiation.

Assessment of non-metallic inclusions That such studies should be in progress is all to the good. Meanwhile it is necessary for production control and for setting acceptance standards between customers and suppliers to have means of evaluating the inclusions content of steel. This has been done traditionally by taking samples of the rolled product from parts of ingots or casts, selected on bases known from experience to give a representative evaluation of cleanness. The bar samples are sectioned, polished and etched, viewed through the optical microscope, and assessed by set procedures, of which there are numerous variations from direct counting and sizing to qualitative comparison with prepared charts showing typical inclusion

The Heterogeneity of Steel

307

distributions. Apart from the high labour content of such procedures, these methods are objectionable on the grounds of the element of judgment required by the observer and the fact that the tedium involved induces fatigue, thus impairing judgment. (There are often even more fundamental objections which I shall discuss later.) These shortcomings have prompted the development in recent years of automatic methods of inclusions assessment designed to perform essentially the same task as the human observer on the same type of samples section but in a much more objective manner. All of these techniques include a scanning principle. Examples of optical scanning instruments are the Quantimet (Image Analysing Computer) of Metals Research Ltd33 and the TI automatic inclusion counter (Cottingham et al.34) developed from a modification by Cottingham of the Vickers projection microscope. In addition to size and numbers of inclusions, some limited information may be obtained about chemical composition based on differences between optical reflectivity of mainly oxide and mainly sulphide inclusions. In electron probe instruments of the type developed by Melford and Whittington= and by Vickers Research Ltd36 substantially more chemical information is obtainable. Some of these instruments include electronic means of further processing signals obtained from the inclusions in order to seek out significant correlations not readily obtainable by the human observer techniques which they aim to replace. There can be no doubt that in relation to the task set for these automatic inclusion assessment instruments they represent a worthwhile advance in technique. With varying levels of sophistication they deliver an objective and speedy description of the inclusion population revealed on a polished and etched section of a sample. We have to realise, however, that the few cm-' of sample examined are but a tiny fraction of the tonnage of steel processed and, furthermore, that what may be observed on a planar section through an inclusion population can be very misleading in relation to the true three-dimensional configuration. The sectioned area of greatly elongated inclusions of the type we showed earlier is very sensitive to the angle of the section with respect to the rolling direction when these are very close, so that the sectioned area, although accurately described, may not be a useful measure of the deleterious effect the inclusion is known to have through its elongation in the third unobserved dimension. Again, in some circumstances the damaging effect of inclusions results from the cooperative action of a number of neighbouring inclusions or inclusion fragments, a circumstance that cannot readily be identified by observation on a single planar section. This effect is strikingly illustrated by some recent observations by Rege et al., 37 of alumina inclusions in low carbon steel. Figure 23a shows a polished section of the as cast steel revealing a cluster of apparently unconnected particles. Observation in the scanning electron microscope of a similar area after deep etching reveals a single alumina inclusion with a highly dendritic form (see Fig. 23b). All of this does not mean that conventional methods of inclusion assessment, whether human or automatic, are useless. On the contrary, they can often give a very good indication of the average quality of steel and provide an important means of continuously monitoring steelworks procedures to ensure adherence to established practices. Thus, although in the course of a year in which some 60 000 tons of En31 bearing steel is

308

Hatfield Memorial Lectures VoL II

Fig. 23

Three-dimensional view of alumina clusters in AI killed low carbon steel. 37 (a) typical alumina clusters at the quarter thickness location of a cross section of the as cast steel, planar view, unetched (x 1500); (b) electron scanning image of an alumnina cluster in a quarter thickness sample of the as cast steel etched for 6 min in 5% solution of bromide in methanol (x 1500).

processed into tube at a particular works, the aggregate area of inclusion noted is only 2 ern? in the course of sample examination over a total area of2000 ern? in about one in six of the incoming casts, this monitoring gives a useful indication of end quality. Even allowing for the similar sampling of every cast before it leaves the steelworks and a ring fracture test on samples of finished tube, it is remarkable that these procedures together ensure adequate control of fatigue performance of the steel and tube for use in its final form as a bearing. Of course, I have not completed the picture because more specialised and infrequent direct fatigue tests by tubemakers on tube rings and by bearing manufacturers on bearings plus knowledge of bearing performance in the field all contribute to the primary selection of a good steelmaking practice. Also, most importantly, continuous

monitoring of this basically good practice includes the tubemaking process itself which,

The Heterogeneity of Steel

309

with its severe demands on metallurgical homogeneity that I have already illustrated, is a very good method of destructive testing. Gross defects arising from relatively rare but large inclusions cause rejection of tube with consequent expense for steel and tubemaker. This involuntary destructive testing gives reasonable protection to the tube customer but the user of plate may be less fortunate in that the processing of ingot to plate is less demanding on steel properties than tube in this particular respect. Thus the problem of lamellar tearing associated with the relatively rare occurrence of larger inclusions will frequently become evident only at a very late stage in the fabrication of complex engineering structures when rectification is very difficult, not to say expensive.

Non-destructive testing Sampling methods of inclusion assessment cannot be expected to provide much, if any, evidence of the presence of the larger infrequent inhomogeneities which are critical to engineering performance. This is the principal reason for the growing interest in recent years in the use of on-line methods of non-destructive testing. Ideally, one requires a method permitting 100% inspection of a product at a speed commensurate with the rate of throughput of the production line, to detect and mark the location of only those defects likely to be deleterious to the performance of the product and all of this at a minimal additional cost. Such defects include, of course, cracks caused by inhomogeneities as well as the original inhomogeneities themselves. I would like briefly to review the progress made towards these aims in the development of ultrasonic techniques for tube inspection in TI. Figure 24 illustrating the flow of ultrasonic energy inj ected into a tube from a single probe shows at once that to achieve 100% inspection of the whole volume with such a system is likely to be an unacceptably laborious operation requiring a relative motion of tube and transducer in a close pitched spiral. This limitation has been overcome in the

Fig. 24

Computer plot of transmission of ultrasonic beam in tube.

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Hatfield Memorial Lectures VoL II

high speed multiprobe ultrasonic tester shown in Fig. 25. In this device there is a circular array of 24 individual probes each covering an arc of 5° of the tube circumference. These are coupled to the tube by water contained in a chamber through which the tube enters and leaves by specially designed seals. The array of probes may be used in a variety of ways to receive echoes from defects within the tube. The sequential transmission and receiving requirements are achieved with electronic switching circuits. Echo signals above a preset level are caused to operate a visual alarm and a paint spray to mark the tube at the appropriate place. The speed of throughput of the tube is related to the desired probability of finding a defect of a given size. In practice axial speeds of the order of 150 ft/rnin are commonly used. Fuller details of the device are to be found in a paper by Kyte and Whittington.Y

Fig. 25

High speed multiprobe

ultrasonic testing equipment

in use for tube inspection.P'

Experience with this high speed tester has laid the basis for a further development with enhanced performance. The new design incorporates the idea of electronic beam swinging commonly used in radar and sonar. The transmitter consists of a small number of separate transducer elements which, if pulsed in sequence with suitably chosen phase delays between each pulse, behave in a cooperative manner to produce a beam of energy in a direction determined by the relative phases of the individual transducers. Figure 26 shows a Schlieren picture of the ultrasonic wavefront produced from a linear array of ten elements from which it can be seen that the beam direction may be changed or the divergence altered as a result of the changes in relative times of pulsing the individual elements. Applying this principle to a circular array of probes around a tube and with the addition of a number of such arrays side by side, it is anticipated that it will be possible to dispense with mechanical rotation of the tube, detect transverse as well as longitudinal defects, and to accommodate tubes of different sizes by simple and quick changes of

electrical controls (Whittington and

COX39).

The Heterogeneity of Steel

311

Fig. 26 Schlieren photographs showing how various wave configurations may be produced by controlling relative phases of a linear array of transmitter elements (parallel to top edge of pictures+"). (a) converging wavefront; (b) diverging wavefront; (c) plane wavefront angled top right; (d) plane wavefront angled to left.

Even with these advanced methods of flaw detection there is still an important practical problem to solve concerning the interpretation of echo signals in terms of the inhomogeneities causing them. For quality control purposes it is customary to compare echo strength with that produced by a standard notch cut into the surface of a tube. If we knew more of the detailed nature of the scattering of ultrasound by well characterised metallurgical inhomogeneities of various types, it is possible that we could employ more sophisticated methods of echo signal analysis permitting better discrimination between harmless and harmful inhomogeneities than is possible by observing signal strength alone. Work is in progress at a number of laboratories on so called ultrasonic spectroscopy to do just this. As yet these studies are not sufficiently advanced to be of any practical value for on-line high speed testing. Another approach to this problem which we have adopted in our Hinxton Hall laboratory is to observe directly the passage of ultrasound in transparent models by use of a sensitive Schlieren apparatus (Marsh and Gunton+'). Pulses of intense ultrasound passing through these models deflect an incident light beam through small angles, typically less than 100 s of arc, and in sufficiently stable and sensitive equipment these pulses appear as bright images in a dark field. An argon jet light source producing 0.1 J.lsflashes at the rate of 240 pps is synchronised to the ultrasonic pulses to within 5 ns, so producing static images as shown in Fig. 27. Here an incident shear wave in a glass model is reflected by a surface containing a defect two wavelengths in length. Beside the

312

Hatfield Memorial Lectures Vol. II

main reflected pulse both shear and longitudinal defect pulses are imaged, and by varying the time delay between light and sound pulse initiation these pulses may be studied as they propagate through the model. Weare only just at the start of these studies, the result shown in Fig. 27 was obtained recently, and it is our hope that they will be a useful guide to possible methods of improved echo discrimination in ultrasonic testing.

Incoming 2MHz shear pulse

Longitudinal pulse from defect

(OO~~

.~fe:jUlses' reflected from

__ Fig. 27

defect

.----'~n/. 2A defect

Schlieren photograph of ultrasonic pulse behaviour in glass model with 2A notch at position indicated.F"

In this lecture I have ranged from the eighteenth century to today, from the ingot to the atom, from the scrap yard to the research laboratory. With my chosen subject I have tried to illustrate the essential historical continuity of scientific enquiry, the essential interplay of researcher, developer, producer and customer, each making his own unique contribution, but always in the light of the needs of all the others, the essential unity of research itself, with no perceptible demarcation between pure and applied nor between one academic discipline and another. If I have succeeded in this, then my task is truly done because these were Hatfield's guiding principles and it is right and proper to recall them in his Memorial Lecture.

The Heterogeneity of Steel

313

REFERENCES 1. W. H. HATFIELD:The Application of Science to the Steel Industry, American Society for Steel Treating, 1928, 88. 2. ]. E. O. MAYNE et al.:J. Chem. Soc., 1950,381. 3. D. A. MELFORD: (a)JISI, 1962, 200, 290; (b)JISI, 1966, 204, 495. 4. R. CASTAING:Thesis, University of Paris, 1951, ONERA publ, no. 55. 5. V. E. COSSLETTand P. DUNCUMB: Nature, 1956, 177,1172. 6. P. DUNCUMB: lSI Conference, Computational Techniques as an Aid in Physical Metallurgy, Leeds, 4-5] an. 1971. 7. C. J. COOKE and P. DUNCUMB: Proceedings of the 5th International Congress on X-ray Optics and Microanalysis, Springer, Tiibingen, 1969,245. 8. M. H. ]ACOBS: Tube Investments Research Laboratories, private communication. 9. E. W. MULLER and J. A. PANITZ: 14th Field Emission Symposium, 1967, Gaithersburg, Maryland, USA. 10. S. S. BRENNER and]. T. McKINNEY: Private communication, to be published in Suif. Sci. 11. P.]. TURNER and M.J. SOUTHON: 'Dynamic mass spectrometry', vol. 1 (eds.) D. Price and ]. E. Williams, Heyden and Sons, London, 1970, 147. 12. J. J. LANDER: Phys. Rev., 1953, 91, 1382. 13. L. A. HARRIS: (a)J. Appl. Phys., 1968,39,1419; (b)]. Appl. Phys., 1968,39,1428. 14. H. L. MARCUS and P. W. PALMBERG:Trans. AIME, 1969,245,1664. 15. R. D. DOHERTY and D. A. MELFORD:]ISI, 1966,204,1131. 16. D. A. MELFORD and D. A. GRANGER: 'The solidification of metals', 289; 1968, London, The Iron and Steel Institute, London, 1968, 289. 17. L. W. PROUD and E. ]. WYNNE: 'Deformation under hot working conditions', The Iron and Steel Institute, London, 1966, 157. 18. H. S. AYERS: Tube Investments Research Laboratories, private communication. 19. H. S. AYERS and T. WILSON: Tube Investments Research Laboratories, private communication. 20. H. S. AYERS and R. SAWLE: Tube Investments Research Laboratories, private communication. 21. Reported in Met. Mat., 1970,4,392. 22. S. BERGH:JISI, 1963,201,510 and 966. 23. P. H. SALMON-COX and]. A. CHARLES:(a)]ISI, 1963,201,863; (b)]ISI, 1965,203,493. 24. F. B. PICKERING: International Conference on Clean Steel, Balatonftired, Hungary, 23-26 June 1970. 25. R. KIESSLINGand N. LANGE: 'Non-metallic inclusions in steel', (Pts I-III), The Iron and Steel Institute, London., 1968. 26. D. M. COTTINGHAM:Deformation Under Hot Working Conditions, The Iron and Steel Institute, London, 1966, 145. 27. C. HOLDEN et al: International conference on clean steel, Balatonftired, Hungary, 23-26 June 1970. 28. W. R. CLARKEel al.: Tube Investments Research Laboratories, private communication. 29. P. J. H. MAUNDER andJ. A. CHARLES:]ISI, 1968,206,705. 30. j A. CHARLES and I. UCHIYAMA:jISI, 1969,207,979.

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Hatfield Memorial Lectures VoL II

31. R. KIESSLINGand H. NORDBERG: International conference on clean steel, Balatonftired, Hungary, 23-26 June 1970. 32. D. BROOKSBANKand K. W. ANDREWS: International Conference on Clean Steel, Balatonftired, Hungary, 23-26 June 1970. 33. C. FISHER: Auto111atic Cleanness AsseSS111entof Steel, The Iron and Steel Institute, London, 1968,24. 34. D. M. COTTINGHAM et al.: Automatic Cleanness Assessment of Steel, The Iron and Steel Institute, London, 1968, 31. 35. D. A. MELFORD and K. R. WHITTINGTON: Proceedings of the 4th International Symposium on X-ray Optics and X-ray Microanalysis, Paris, 1965; 1967, Paris, Hermann. 36. (a) R. BEADLEet al.: Proc. lEE, 1969, 116, 6; (b) K. RIDAL: Automatic Cleanness Assessment of Steel, The Iron and Steel Institute, London, 1968, 40. 37. R. A. REGE et al.: Met. Trans., 1970, 1,2652. 38. G. H. KYTE and K. R. WHITTINGTON: Industrial Ultrasonics, Proceedings of Lough borough Conference on Ultrasonics, Sept. 1969, 145. 39. K. R. WHITTINGTON and B. D. Cox: Ultrasonics, 1969,7,20. 40. K. R. WHITTINGTON and B. D. Cox: Tube Investments Research Laboratories, private communication. 41. J. H. GUNTON and D. M. MARSH: Industrial Ultrasonics, Proceedings of Loughborough Conference on Ultrasonics, Sept. 1969,257. 42. D. M. MARSH and V. M. BABOROVSKY:Tube Investments Research Laboratories, private communication.

TWENTY-NINTH

HATFIELD

MEMORIAL

LECTURE

Ferrite R. W. K. Honeycombe At the time the lecture was given Professor R. W. K. Honeycombe was Goldsmith's Professor if Metallurgy, University of Cambridge. The lecture was presented at Firth Hall, University of Sheffield, on 6 December 1979.

The morphology of ferrite grains formed from austenite is first considered, then the refinement of ferrite grain size by controlled rolling as exemplified by HSLA steels. The influence of alloying elements is briefly examined, followed by a detailed discussion of the precipitation of carbides during and after the formation of ferrite. It is shown that the yI a interfaces are preferred sites for the precipitation of alloy carbides (interphase precipitation)",which occurs on both planar and curved boundaries. In the former, the precipitation is associated with the movement of small ledges across the yla interface, whereas the curved bands of precipitate are formed on high energy boundaries which move by bowing around the particles. A change in the nature of the yIa interface often leads to fibrous carbide growth, while more rapid rates of cooling lead to the formation of supersaturated ferrite and subsequent precipitation of carbides on dislocations. The dislocations in the ferrite are shown to play an important role in the subsequent coarsening of the carbide dispersion, whether they form at the interfaces or on the dislocations. Finally, the observations on the formation and growth of carbide dispersions are related to the structure and properties of micro-alloyed steels.

Dr William Hatfield was one of the outstanding men bred in the smoke and fire of the Sheffield steel industry in the first half of this century. The smoke is gone, but the fire remains and the great store of scientific and technological knowledge built up by pioneers such as Hatfield has been added to by successive generations of metallurgists and engineers dedicated to keeping Sheffield in its prime position as one of the major alloy steel manufacturing centres of the world. As part of this evolutionary process, scientific and technical education in Sheffield has more than done its share, starting with the foundation of Firth College, the centenary of which is being celebrated this year. The study of engineering and metallurgy in Sheffield was already well established in the latter half of the nineteenth century at the Sheffield Technical School. A great many of the early Sheffield steel men received their education at these institutions which joined with the Sheffield Medical School to become University College in 1897. This was where Hatfield himself studied under Professor J. O. Arnold and received the Mappin Medal in 1902. In 1905 University College became the University of Sheffield, which today from

315

316

Hatfield Memorial Lectures VoL II

its applied and pure science departments makes a major contribution in manpower and research not only to the steel industry in Sheffield, but to the whole of the industry in the United Kingdom. It is now truer than ever that steels are the most important group of engineering materials, for they are continually evolving to meet new needs and challenges, and this is really the main justification for doing research in this field. My topic tonight is concerned with the low temperature form of iron, generically referred to as ferrite, the dominant constituent in very many irons and plain carbon and alloy steels. Research work in this area is important because increasing tonnages of low alloy steels are being made in which the route to reliable mechanical properties is directly from austenite to ferrite, without the need for quenching to form martensite which then requires tempering.

TRANSFORMATION

FROM AUSTENITE TO FERRITE

In pure iron under equilibrium conditions the transformation from fcc austenite to bee ferrite occurs at about 910°C. While rapid cooling can alter the transformation temperature, the addition of alloying elements causes transformation to occur over a wide temperature range and allows a closer investigation of the microstructural changes involved. I shall first illustrate these changes by reference to the simple Fe-C system choosing a steel with O.32%C, which has been heat treated to allow the ferrite reaction to commence, but in which the reaction has been interrupted by quenching. Figure 1 shows some of the microstructures encountered. In the higher temperature range (7S0-850°C), the y ---7 a reaction is nucleated predominantly at the austenite grain boundaries as grain boundary allotriomorphs which are equiaxed or lenticular in form (Fig. 1a). Intergranular crystals are also nucleated (Fig. 1a) and if the reaction is allowed to go to completion, an equiaxed polyhedral grain morphology is established. As the transformation temperature is lowered, the allotriomorphs develop planar facets on at least one side and often on both sides of the boundary (Fig. 1b), indicating that low energy crystallographic interfaces are beginning to predominate instead of randomly oriented curved high energy interfaces." The ultimate result of this trend is the development of a well defined lath shaped ferrite which has a precise crystallographic relationship with the austenite in which it forms. This Widmanstatten ferrite is readily recognised in the optical microscope (Fig. 1c, d), and becomes more prominent as the y ---7 a transformation temperature is depressed below 7SO°C. The ferrite morphologies so far described arise by nucleation at preferred sites such as austenite grain boundaries, particles of second phases or inclusions, and subsequently by thermally activated growth processes. A very large number of y/ a interfaces are planar and of relatively low surface energy. These can be observed dynamically during the transformation by use of the photoemission electron microscope (Fig. 2), which confirms the predonlinance of planar boundaries.? and shows directly that these boundaries grow by the lateral movement of small steps. These steps vary in height from several micrometres to a few atomic distances, depending on the composition, temperature and other

Ferrite

317

Fig. 1 Optical micrographs ofFe-O.34C alloy showing particle isothermal transformation in the range 800-700°C (Ricks): (a) grain boundary allotriomorphs; (b) side growths of ferrite; (c) Widmanstatten ferrite; (d) transgranular Widrnanstatten ferrite.

variables. They occur over the whole transformation range, participating in the formation of equiaxed ferrite, but become the dominant mechanism in the growth of Widman statten ferrite, where the semi-coherent crystallographic interfaces between 'Y and a are not mo bile? (Fig. 3). At still lower temperatures, ferrite changes its morphology yet again to adopt a blocky type of structure which is now referred to as bainitic fcrrite.i'-" This ferrite is not only found associated with iron carbide in iron-carbon alloys, but also in carbon free alloys where the phase is precipitate free ferrite. Figure 4 shows the occurrence of this type of ferrite in an Fe-l0%Cr alloy transformed at 525°C; it is difficult to distinguish from Widmanstatrcn ferrite in the optical microscope, but does exhibit a more blocky habit. However, a much clearer distinction is found using electron microscopy, which allows us to distinguish further between these various types of ferrite by revealing the dislocation structures within the grains, and also by permitting the detailed examination of subgrain

structures. The dislocations arise as a result of the volume change which takes place when

318

Hatfield Memorial Lectures VoL II

Fig. 2

Series of photo emission electron micrographs of same area of Fe-O .19C-ll. 7Cr during transformation at 650°C; new ferrite grain growing into three austenite grains with interfaces <Xl' <X2 and <X3 (Edmonds and Honeycombe).

austenite transforms to ferrite (-1 % in pure iron at the equilibrium transformation temperature). This leads to the development of internal stresses and their subsequent reliefby dislocation generation and movement. Clearly, the lower the temperature at which ferrite is allowed to form, the larger will be the volume change and the greater the dislocation density which results. Figure Sa and b contrasts the dislocation structure in equiaxed ferrite formed at 800°C, and Widmanstatten ferrite formed at 650°C in an Fe10%Cr alloy.> The subgrain structure of equiaxed ferrite usually comprises low angle boundaries (e.g. Fig. Sa), which form as a result of dislocation climb at the higher

transformation temperatures. In contrast, the Widmanstatten

laths often contain a

320

Hatfield Memorial Lectures VoL II

Fig.5 TEM ofFe-l0Cr; isothermal transformation between 800 and 525°C (Bee): (a) 5 min at 800°C, equiaxed ferrite; (b) 60 min at 675°C, Widmanstatten ferrite; (c) 3 min at 525°C, bainitic ferrite.

FERRITE GRAIN SIZE The familiar Hall-Petch relationship between the yield stress and grain size (Fig. 6) shothe strengthening to be achieved by attaining fine ferrite grain size in steels. Twenty-fi years ago, mild steels would be expected to have grain sizes in the 20-50 ~m range al yield strengths between 200 and 300 MN m='. However, the development of micr alloyed steels in the intervening period has led to ferrite grain sizes typically in the ran 5-10 ~m and yield strengths between 450 and 550 MN rn='. This dramatic improveme in mechanical properties results to a large degree from the grain size refinement which l: been achieved by the addition of small concentrations (> 0.1 wt-%) of strong carbi

320

Hatfield Memorial Lectures VoL II

Fig. 5 TEM ofFe-l0Cr; isothermal transformation between 800 and 525°C (Bee): (a) 5 min at 800°C, equiaxed ferrite; (b) 60 min at 675°C, Widmanstattcn ferrite; (c) 3 min at 525°C, bainitic ferrite.

FERRITE GRAIN SIZE The familiar Hall-Petch relationship between the yield stress and grain size (Fig. 6) shows the strengthening to be achieved by attaining fine ferrite grain size in steels. Twenty-five years ago, mild steels would be expected to have grain sizes in the 20-50 ~m range and yield strengths between 200 and 300 MN m-2• However, the development of microalloyed steels in the intervening period has led to ferrite grain sizes typically in the range 5-10 ~m and yield strengths between 450 and 550 MN m-2. This dramatic improvement in mechanical properties results to a large degree from the grain size refinement which has been achieved by the addition of small concentrations (> 0.1 wt-%) of strong carbide

Ferrite

321

forming elements such as Nb, Ti and V to essentially mild steels, and by the use of controlled rolling for fabrication.v-? The grain refinement arises primarily as a result of the pinning of austenitic grain boundaries and of dislocation arrays within the deformed grains, which delays the recrystallisation of austenite until it has been more heavily deformed. In other words, the critical strain required for the onset of recrystallisation is substantially increased so that, when it does take place, there is a higher rate of nucleation, and thus a finer austenitic grain size results. Pinning of dislocations in the austenite of a low alloy steel is illustrated in Fig. 7.

Fig. 6

Fig. 7

Effect of ferrite grain size on the yield strength of O.09C-l.SMn steel and a similar steel with O.05%Nb added (Le Bon and de Saint-Martin).

Dark field TEM (VC reflection) Fe-1V-O.2C aged at 875°C prior to transformation at 700°C; precipitation ofVC on austenite sub-boundaries and dislocations (Walker).

322

Hatfield Memorial Lectures VoL II

The austenite formed following recrystallisation is further deformed during rolling until the y ~ a transformation takes place. The plastic deformation has two effects on the transformation: (i) the y ~ a reaction is accelerated (ii) the resulting a grain size is further refined The first point is demonstrated by some results obtained by Walker" on En24 (lCr1.5Ni-O.25Mo-O.36C), which was deformed in the metastable austenitic condition at 550°C then transformed at 650°C. Figure 8 shows that increasing plastic deformation substantially accelerates the y ~ a transformation. It seems clear that the deformation has increased the rates of nucleation and growth of ferrite in austenite; this is confirmed by microscopic examination which reveals that ferrite nuclei not only occur at the austenite grain boundaries, but also within the grains, presumably at dislocation concentrations or sub-boundaries. Figure 9 gives an example of intra granular nucleation of ferrite in deformed austenite. The encouragement of intragranular nucleation of ferrite in austenite

0:z 100

Q 80

~ z ~

50

LL

40

0 (/)

Ni-Cr-Mo steel at 550°C

z 20 4:

~ •....

Fig. 8

Fig.9

0 102

103 TIME,s

104

Effect of plastic deformation at 550°C on the rate of transformation of austenite in En24 at 650°C (Walker).

Optical micrograph ofEn24 1.35Ni-1.07Cr-0.27Mo-0.37C deformed 10% at 550°C prior to transformation at 650°C; intra granular idiomorphs of ferrite (Walker).

Ferrite

323

is an important aspect of the achievement of fine grain sizes, which could well repay further attention. Not only may the finely divided phases in the austenite inhibit the recrystallisation of austenite, but they may also act as nuclei for ferrite during the phase transformation. Application of the above principles to micro alloyed steels has resulted in ferrite grain sizes in the range 5-10 J..lmbeing regularly achieved." Atypical result? is shown in Fig. 6 in which the introduction of niobium has substantially extended the range of grain sizes which can be achieved by controlled rolling of carbon-manganese steels.

EFFECT OF ALLOYING ELEMENTS ON TRANSFORMATION KINETICS In pure iron the 'Y~ a reaction is extremely rapid but, as is well known, the addition of alloying elements both in interstitial and substitutional solid solution retards the rate of transformation. Carbon has a large effect because of its much higher solubility in austenite than in ferrite. Consequently during the "{~ a transformation carbon is precipitated as carbide, and the rate controlling process in Fe-C alloys will be the diffusion of carbon in austenite. Metallic alloying elements are also very effective. Figure 10 shows the TTT curve for an Fe-l0%Cr carbon free alloy, in which the reaction exhibits the familiar C curve behaviour but occurs more slowly than in pure iron.> If sufficient alloying element is added, this permits the steel to bypass the ferrite reaction and to transform at lower temperatures to martensite by a diffusionless type of transformation. While convincing thermodynamic arguments have been advanced to explain this behaviour, I think it is fair to say that the distribution of metallic alloying elements between the "{, U, and carbide phases during the transformation of steels has been largely a matter of speculation. The much greater mobility of the interstitial solutes, carbon and nitrogen, ensures that at higher temperatures, e.g. 850-700°C, partition of carbon between austenite and ferrite

a

equlaxed

------ws <,

equlaxed and Widmanstatt(ln

a

-----Bs

102 LOG TIME,

Fig. 10

s

103

TTT diagram showing critical temperatures for Widmanstatten ferrite, bainite and martensite reaction in Pe--I OCr (Bee).

324

Hatfield Memorial Lectures VoL II

takes place; this continues at lower temperatures, e.g. when bainitic ferrite is formed, but increasingly the ferrite is likely to become supersaturated with respect to carbon. Turning to the substitutional solutes, the usual metallic alloying elements in steel do partition between austenite and ferrite under equilibrium conditions, and recently quantitative techniques such as scanning transmission electron microscopy (STEM) have become available which allow the different regions of a microstructure to be observed and analysed on a very fine scale with a typical resolution of 50· nm. To take a good example, the y and a iron can coexist in equilibrium in a duplex stainless steel (e.g. containing 26Cr-5Ni-1.3Mo-0.03%C), as shown in Fig. 11a, and STEM analysis shows clearly how the chromium has distributed between the ferrite and the austenite (Fig. 11b) as predicted by the equilibrium diagram. 10 However, turning to fully transformable steels,

+:

,

'++++'+ ! I J I

farrita

: austenlte •I

+

farrita

+

+ +i

! + ++ i , I

23~

o

I I

I

:

~ __ ~'~

~t

2

DISTANCE,

~~~

3

4

urn

Fig. 11 Local chemical analyses using scanning transmission microscopy and an energy dispersive technique: (a) duplex stainless steel; (b) typical Cr scan across similar microstructure (Southwick).

Ferrite

325

partial transformation does not represent equilibrium, and only when 100% ferrite is achieved is this reached. In these circumstances, partition of the metallic alloying elements between ferrite and austenite does not necessarily occur, yet we know that these elements often markedly retard the 'Y ~ ex reaction. STEM analyses have been recently carried out by my colleagues Drs P. R. Howell and R. A. Ricks-! on a 10%Cr-0.2%C steel, a 1%V-0.2%C steel, and on Fe-Cu-Ni alloys. In no case, and at no transformation temperature between 800 and 600°C, was partitioning of metallic elements detected, the limit of resolution being about 50 nm. Figure 12a shows a typical traverse across the microstructure of a partially transformed 10%Cr steel in which there is a uniform distribution of chromium in both transformed and untransformed regions. This implies that volume diffusion of chromium is not taking place during the transformation at least

Fig. 12

STEM analyses at yla interface in a partially transformed 10%Cr steel: (a) typical microstructure at 700°C; (b) Cr concentration scan (Ricks).

326

Hatfield Memorial Lectures Vol. II

within the limits of detection (- 50 nm). However, the chromium concentration at the interface is sometimes found to be significantly higher (Fig. 12b), which is consistent with chromium diffusion within the interface between yand a. Dilatometric studies of the 'Y ~ a transformation in several alloy steels, particularly microalloyed steels based on 0.25%V, and on carbon free iron-nickel-copper alloys, has revealed well defined discontinuities in the isothermal transformation curves12,13 (Fig. 13). These occur well below the AC1 and AC3 temperatures, and it is perhaps tempting to associate them with a change in the distribution of the alloying elenlent, with partition to austenite occurring above the discontinuity but no partition below. However, the above results from STEM analysis tend to eliminate this, and point to the diffusivity of the alloying elements being limited to the y/a interface. This then supports the theory of 'solute drag' for explaining the retardation of the y ~ a transformation by metallic alloying elements, first put forward by Lucke and Detert.I" They proposed that moving interfaces collected solute atoms during the transformation and these atoms then exerted a drag on the interface, the movement of which would then be controlled by the interfacial diffusivity of the solute atoms. As we shall see later, this solute movement at the interphase boundary is particularly important when carbides are formed at the same time as the ferrite.

10

TIME,s

100

1000

Fig. 13 TTT curves for Fe-O.26V-O.022N-O.020C obtained from high speed dilatometer: 5 and 95% transformation (Balliger).

CARBIDE PRECIPITATION DURING TRANSFORMATION In plain carbon and in many alloy steels the y ~ a transformation is accompanied by the precipitation of carbon as iron carbide, most familiar in pearlite, an important constituent of many engineering steels. However, the introduction into the steel of stronger carbideforming elements than iron (Fig. 14), e.g. Cr, Mo, W, Nb, Ti and V, means that these elements will combine preferentially with the carbon, and consequently alloy carbides would be expected to form during the 'Y ~ a transformation. While in some cases these elements

Ferrite 327 can lead to the formation of alloy pearlites analogous. to cementitic pearlite (Pig. 15), .oficnihe morphologies observed are dramatically different from that. of pearlite. 1~ .This is true of wide range of alloy steels, but the most significant results are obtained when using very strong carbide formers such as Nb, Ti and V,. where additions to the steel of less than 0.1 wt""-%result in alloy carbide precipitation during and after the y ~ ex transformation.

a

TiB2 NbB2

ZrB2

T01B2

HfB2 .. 160

0

40

80

120

Cr23C6 TiC

VCx

Cr7C3

Fe3C

(Cr

C ) 3 2

Mo2C

NbC

ZrC

(MoC)

(Nb C) 2 HfC

W2C

raC

(we)

CT012C)

160 TiN

160

0

40 Cr2N

(Mo2N)?

NbN

ZrN

HfN

80

120 VN

TaN

80

120 -b.H298

40

0

K , kJ 19 atom

Fig. 14 Enthalpies of formation of carbides, nitrides and borides (from K. L. Schick: Thermodynamics of Certain Refractory Compounds, Academic Press, New York, NY, 1966).

Fig.15

SEM of Fe-10Cr-O.2C transformed 1 h at 700°C; MZ3C6 pearlite (Howell).

328

Hatfield Memorial Lectures VoL II

This brings us back to HSLA or microalloyed steels, in which we have already examined the effect of grain size refinement. Looking again at the comparison of a C-Mn steel with a Nb containing steel (Fig. 6), it can be seen that the yield stress versus grain size curve for the latter is displaced to higher strength levels. This important fact, first emphasisedby Morrison and Woodhead, 16 implies that there is, in addition to grain refinement, another strengthening mechanism associated with the presence ofNb, namely dispersion strengthening due to carbide precipitation. The solubilities of NbC, VC and TiC in austenite are at least an order of magnitude greater than in ferrite. (Fig. 16), so that precipitation during the transformation from austenite to ferrite can be expected. This precipitation has been studied in detail in low alloy steelswith Nb, Ti and V, as well as in higher alloy steels containing Cr, Mo and W,17 and while the morphologies are often complex it is now possible to categorise them and offer explanations on how they arise. At least four different morphologies of carbide have been established:

",

"" , ,, "-

"-

,

" , "-

"",

""- ,

", ""

VC in

terrtte

Fig. 16

1. 2. 3. 4.

Solubility products of VC and NbC in austenite (Aronsson). Dashed line: VC in ferrite.

Interphase precipitation (planar) Interphase precipitation (curved) Carbide fibre growth Precipitation from supersaturated ferrite

These shall be discussed briefly in tum.

Ferrite

329

Interphase Precipitation (Planar) Perhaps the most interesting of the precipitate distributions observed is a characteristically banded structure in which the usually very fine carbide particles are arranged in regularly spaced bands (Fig. 17). The bands have been shown to be parallel to the 'Y/ a interface, which suggests that nucleation of the particles has taken place periodically at the boundary between the 'Y and a phases, hence the term interphase precipitation.P' Much of the basic work which has led to an understanding of the mechanism of this phenomenon has been carried out on high chromium steels,19,20 in which it occurs on a much coarser scale. than in microalloyed steels. Figure 18 shows M23 C6 interphase precipitation in a partially transformed 12%Cr steel, where the 'Y/ a interface is shown to contain large steps (- 1 urn). It is obvious that the precipitate band spacings are related to the step height. Detailed electron microscopy has provided the following information: (i) (ii)

the bands of precipitate are closely parallel to the 'Y/ a interface the interface grows by the sideways migration of steps

Fig.17

TEM ofFe-0.75V-0.15C

transformed 5 min at 725°C (Batte and Honeycombe).

Fig. 18

TEM of Fe-12Cr-0.2C isothermally transformed 30 min at 650°C; ferrite and MZ3C6, coarse stepped interface (Campbell and Honeycombe).

330 (iii) (iv)

Hatfield Memorial Lectures VoL II the precipitate nucleates on the non-migrating planar interface, and not on the moving steps only one variant of the precipitate orientation relation is exhibited in a particular region of the interface

More .rccently.e" using similar steels, but by a detailed study of the crystallographic relationships between the ferrite, M23C6, and austenite retained by quenching, Howell and Bee have shown that the interphase precipitation of M23C6 is related both to the ferrite and to the austenite: (111)1'1 1(011)al 1(111)M23c6 [~01]1'1 1[111 ]al 1(1 01)M23C6 Such observations imply that the interphase precipitate bands are associated with the coherent growth of the ferrite related to the austenite by the Kurdjumov-Sachs relationship, and that the carbide is nucleated on the austenite/ferrite interface and not within one of the phases.

Interphase Precipitation (Curved) The very straight bands of precipitate are undoubtedly associated with low energy planar yla interfaces between orientation-related austenite and ferrite. However, there is now overwhelming evidence that in many cases the bands are curved.l=!? and sometimes severely distorted implying that precipitation can also take place at high energy curved yla boundaries. Figure 19a gives an example of such a boundary in a 12%Cr steel. It is assumed that high energy curved yla boundaries of this type do not migrate by step propagation, but that the boundary escapes locally from precipitate particles by bending between them (Fig. 19b). This means that yla boundaries between unrelated grains of austenite and ferrite can also create bands of precipitate, and explains why interphase

Fig. 19 TEM ofFe-12Cr-0.2C: (a) isothermally transformed 60 min at 625°C, irregular ,,{Ia interface (Campbell and Honeycombe); (b) isothermally transformed at 655°C, bowing of ria interface away from MZ3C6 particles (Ricks). Arrow indicates "(Ia interface.

Ferrite

331

precipitation can be quite dominant in some microstructures, without the need to postulate that only orientation related "{/ a boundaries moving by step migration take part in the transformation. This latter view is not valid, particularly at high transformation temperatures, when there is much movement of high energy interphase boundaries. The existence of the two mechanisms of interphase precipitation has been recently confirmed by Ricks.s! who examined the "{~ a transformation in Pe--Cu+Ni alloys which can precipitate £-copper during the transformation. Figure 20 shows the types of transformation structures obtainable at 700-720°C; in the optical micrograph the equiaxed ferrite is apparently featureless (Fig. 20a), but thin foil electron microscopy shows both planar and curved bands of £-Cu precipitate (b-d) making it quite evident that both coherent orientation-related ,,{Ia boundaries and high energy curved boundaries were responsible for the copper dispersion within the ferrite.

Fig. 20 Fe-2Cu-2Ni showing E-Cu precipitate distribution (Ricks): (a) 5 min at 700°C (optical); (b) 20 min at 720°C (TEM); (c) 5 min at 700°C (TEM); (d) 5 min at 720°C (TEM).

Growth of Fibrous Carbides in Ferrite In many alloy steels growth of fine fibrous carbide is observed starting from former austenite grain boundaries,22,23 The fibres are usually between 10 and 50 nm diameter and grow in a direction approximately normal to the boundary (Fig. 21a). They are often

332

Hatfield Memorial Lectures Vol. II

Fig. 21 Fibrous carbide morphology: (a) TEM ofFe-4Mo-O.23C transformed 30 min at 700°C (Berry); (b) EM extraction replica of Fe-1V-0.2C cooled 50 K mirr ' from 1150°C (Edmonds). Arrow indicates y grain boundary. observed in close association with interphase precipitation, and frequently within the same ferrite grain, although not necessarily orginating at the same interface.F" There is, however, much evidence to suggest that even at a particular ,,{Ia interface the morphology can change from interphase precipitation to fibre growth, then back again23 (Fig.

21b). The conditions which favour fibre growth appear to be those which bring the alloy closer to equilibrium, for example, high transformation temperatures of addition of elements which slow down the "{ ----7 a reaction, e.g. nickel and manganese. These variables might be expected to encourage the growth of high energy incoherent ,,{Ia interfaces, with which fibre growth appears to be associated. Recent work by Howell et al.20 on high chromium steels has confirmed this view, and demonstrated that the ferrite and associated carbide fibres are not orientation related to the austenite grain in which they are growing. This is observed often at a constant transformation temperature, but can be stimulated. by changing the temperature during transformation.P? The ,,{Ia +

Ferrite

333

carbide boundaries are usually curved (although at high magnifications they may appear straight), and do not move by step propagation, a characteristic they share with the more familiar pearlite nodules growing in austenite. A typical interface between the phases is shown in Fig. 12a for a 10%Cr-O.2%C steel partly transformed at 700°C; the fibrous carbide in this case is M23 C6 which is unrelated to the austenite ahead of the interface (c£ interphase precipitation of M23C6); the y/a interface is free from steps. It seems clear that interphase precipitation and carbide fibres are competing carbide morphologies in ferrite. On the one hand interphase precipitation can occur on low and on high energy boundaries with or without the migration of steps along the interface. This morphology occurs over a wide transformation temperature range, and is particularly prominent when the transformation rate is high. In contrast, fibrous carbide growth is only associated with higher energy y/a interfaces, and does not occur when steps migrate along the interface. One difficult question is how interphase precipitate is replaced by fibre growth, sometimes temporarily on the same transformation front (Fig. 21b). In these circumstances it is probable that the mode of growth of the y/a interface changes from step migration to overall boundary migration. This could be achieved by the slowing down of step movement to the point where nucleation of carbides occurs on the steps and immobilises them. An example of this type of precipitation at steps is shown in Fig. 22, in which M23C6 has precipitated on the ledge which previously has generated interphase precipitate.

Fig.22

M23C6 centred dark field EM ofFe-l0Cr-O.2C: yla boundary with precipitation ofM23C6 on a step (Ricks).

Precipitation from Supersaturated Ferrite Systematic examination offerrite in experimental Fe-O.2SV-O.04(C, N) alloys= (related to microalloyed steels), and in Pe--Ni+Cu alloys13,21 has shown that at lower transformation temperatures, precipitation does not necessarily occur as the transformation takes place, but may occur at a later stage within the ferrite. In these circumstances, the

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Hatfield Memorial Lectures VoL II

nucleating sites within the ferrite are predominantly dislocations, which arise as a result of the volume change during the transformation. This type of precipitation is illustrated in Fig. 23 for an Fe-2Ni-2Cu alloy transformed at 700°C. It can occur in equiaxed ferrite side by side with interphase precipitation, but becomes more significant as the transformation temperature is lowered and Widmanstatten ferrite predominates in the microstructure. Nevertheless interphase precipitation, for example of £-Cu has been observed in association with Widmanstatten ferrite at transformation temperatures as low as 600°C, under which conditions there is also extensive precipitation on dislocations (Fig. 24). The origin of these complex microstructures can be better understood if it is borne in mind that during an isothermal transformation ferrite can nucleate within a few seconds,

Fig. 23

TEM ofFe-2Cu-2Ni isothermally transformed 20 min at 700°C; precipitation of E-CU from supersaturated ferrite, mainly on dislocations (Ricks).

Fig.24 TEM ofFe-2Cu-5Ni isothermally transformed 10 min at 590°C; Windmastatten ferrite plate with precipitation of e-Cu on dislocations followed by interphase precipitation (Ricks).

Ferrite

335

but the transformation may not be complete for several minutes depending on the temperature and composition of the alloy. The first formed ferrite may well be supersaturated and precipitate later on dislocations, whereas the ferrite which forms later may exhibit interphase precipitation even within the same grain (Fig. 24). Similar considerations apply to alloys transforming during continuous cooling, where the ferrite and precipitate can nucleate and grow over a wide temperature range. The last possibility is that ferrite will form then remain supersaturated to room temperature. This is particularly likely to happen in alloys transformed at very low temperatures «600°C), or in alloys which are rapidly cooled through the transformation range. At low transformation temperatures, the driving force of the y ~ a reaction is so great that it overcomes any solute drag effect which may occur at the interphase boundary. Moreover the mode of growth of ferrite changes from a diffusion controlled situation to a mechanism involving shear with the result that, when bainitic ferrite is formed, it is invariably supersaturated. This has been found for an Fe-7Ni-2Cu alloy which transforms to bainitic ferrite in the range 480-525°C.

STABILITY OF DISPERSIONS IN FERRITE The carbides of niobium, titanium and vanadium can exist as extremely fine dispersions in ferrite. However, there is now strong evidence that the initial dispersion formed during or just after the y ~ a transformation can rapidly change in character at the transformation temperature as a result of coarsening. In systematic work on simple laboratory steels based on microalloyed steels containing vanadium and titanium, Dunlop-? and Balliger-? have not only determined the mechanism by which growth of the dispersions occurs but have also established ways in which the coarsening can be controlled. The most striking process is the coarsening of the banded interphase precipitation ofVC and TiC (and by implication NbC). The steels used had low carbon contents (0.04-0.1%), and at 790°C the y ~ a transformation was complete in about 100 s; however, even after only 135 s at 790°C, the banded precipitate of V(C . N) already exhibited a bimodal distribution with coarsening of apparently random particles (Fig. 25a and b). Detailed structural investigation showed that the particles which were coarsening all lay on the random dislocation network which is usual in these structures, and this was confirmed using specimens deformed deliberately prior to coarsening treatment (Fig. 25c and d). This role of dislocations in the coarsening process was confirmed by analysis of the coarsening kinetics which, early on in the reaction, showed a dependence of mean particle size on (time)1/2, but for most of the coarsening the particle size varied as (time)1/5 (Fig. 26), which would be expected if dislocation pipe diffusion were the rate controlling process in coarsening. These results also emphasise the need to look at the microstructures after very short aging times, if the true mechanism occurring during the 'Y ~ a transformation is to be studied, rather than

subsequent coarsening phenomena.

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Hatfield Memorial Lectures VoL II

Fig. 25 Fe-O.26V-O.02C-0.022N: (a) and (b) transformed for 135 s at 790°C showing bimodal distribution ofV(CN) with (a) bands visible; (b) different area and tilt; (c) and (d) 5% cold deformation and aged 25 h at 740°C, all V(CN) particles now lie on dislocations (c) bright field EM; (d) same area as (c) but weak beam EM (Ballinger).

E c

3

LOG TIME, s

Fig.26

4

5

Coarsening behaviour of Fe-O.26V-0.02C-O.022N at 790 and 740°C; semi-log plot of mean particle size versus time (Balliger).

Ferrite

337

Some further observations by Dunlop and by Balliger on the coarsening behaviour are of particular practical concern because they indicate ways in which the initial structures can be refined and the coarsening slowed down. For example, the coarsening of vanadium carbide in ferrite can be substantially reduced by replacing a proportion of the vanadium atoms in VC by titanium atoms; both VC and TiC are fcc and exhibit mutual solubility. The resulting mixed carbide not only exhibits a finer initial interphase precipitation, but also coarsens more slowly than either VC or TiC in ferrite. In similar vein, Balliger has replaced part and all of the carbon in vanadium carbide by nitrogen without changing the overall stoichoimetry of an Fe-O.2SV-O.04(CN) alloy, and has found that the nitrogen containing dispersion coarsened about fifty times more slowly than the straight vanadium carbide dispersion (Fig. 27). This substantial effect may result partly from the reduction in solubility of the precipitate (VC is much more soluble in ferrite than is VN), which would adversely affect the early stages of coarsening where a t1/2law applies (Fig. 26), and when the reaction appears to be interface.controlled.

Fig. 27

Comparison of coarsening behaviour at 790°C ofFeVC (Fe-O.27V-O.OSC) with that of FeVCN (Fe-O.26V-O.02C-O.022N) (Balliger).

DIRECT STUDY yla INTERFACE It should now be clear that the microstructure of ferrite in alloy steels is to a large extent

determined by the nature of the y/ a interface, but this interface is elusive if only because it is difficult to retain ferrite and austenite together at room temperature. However, by choosing suitable high alloy steels it is possible not only to retain the y and a phases at room temperature, but also to study in detail the structure ofy/a interfaces formed over a wide range of temperature by varying the heat treatment. Howell and Southwick-" have used a 26Cr-6Ni-l.3Mo low carbon duplex stainless steel for this purpose, and have studied y/a interfaces formed in the temperature range 1000-300°C. Above about 850°C austenite growing in ferrite (the reverse situation to that in low alloy steels) has a

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Hatfield Memorial Lectures VoL II

coarse rod-like morphology; planar interfaces corresponding to (lll)y/ /(110)a which contain no intrinsic dislocations are observed (Fig. 28). Other boundaries have one set of observed dislocations, the spacing of which depends on the deviation from the 'coherent' boundary orientation. These dislocations can' be explained in terms of the 0 lattice theory of intercrystalline interfaces. Ledges are infrequent in these high temperature boundaries, but when the transformation temperature is below 750°C, many more steps or ledges are observed. A number of these steps are shown in a planar y/a interface in Fig. 29.

Fig.28

Duplex stainlesssteel (26Cr-Sni-l.3Mo-O.03C); "{formed in a at 900°C (Southwick). Planar coherent interface arrowed.

Fig. 29

Dark field micrograph of duplex stainless steel; steps formed on a planar "{Ia interface (Southwick).

In this steel M23 C6 can be precipitated at the interphase boundary, so it is possible to observe interactions of the precipitate with the y/a boundary.F? Where there is a well defined Widmanstatcen boundary, the growth is predominantly by step migration; however, the movement of random interphase boundaries is more complex. It has been found that movement can still be by ledge migration, but growth by bowing between precipitate particles is also important. Both these mechanisms can be seen in Fig. 30, the arrows indicating the points where a bowing mechanism is in progress. So work on the model alloy has confirmed that the two main mechanisms ofy/a boundary movement

actually occur, and are found sometimes side by side on the same boundary.

Ferrite

339

Fig. 30 Duplex stainless steel showing y/a interface with precipitate M23C6 arrowed; both ledge movement and bowing of the boundary are taking place (Southwick).

STRUCTURE OF MICROALLOYED STEELS During this lecture I have from time to time referred to microalloyed steels as illustrating several of the phenomena I have dealt with. They are, I believe, the most outstanding example of the direct exploitation of the y ~ a transformation. Steels richer in alloying elements are usually quenched and tempered, while microalloyed steels are generally brought to their final condition by controlled cooling after hot rolling. Whereas twentyfive years ago, ferrite was the weaker (but tougher) partner which, in combination with pearlite in different proportions, made up the maj ority of structural and engineering steels, today it is often the dominant microstructural component in a range of steels that make little use of pearlite. To achieve this, the carbon content has been lowered, in many cases below 0.1 wt-%, giving obvious advantages for welding and general fabricability. At the same time, the strength has been more than doubled in comparison with earlier ferrite/pearlite steels. This has been achieved in two ways: (i) by ferrite grain size refinement (ii) by dispersion strengthening of the ferrite I have tried to show how small concentrations of strong carbide forming elements added to the steels achieve both of these objectives. Firstly, during the hot rolling of the steels, fine carbide particles form in the austenite and control its recrystallisation behaviour, and through this the final ferrite grain size is markedly refined. Secondly, carbides are precipitated during and immediately after the y ~ a transformation in several ways which lead to a substantial strengthening of the ferrite. The combined effect of these phenomena leads to steels with yield strengths in the range 400-600 MN m-2 and very good toughness. There seems no reason why even stronger steels of this type should not be made. I have only had time to touch on a few important aspects of ferrite in this lecture, and I have tended to concentrate on work done in my own laboratory, fully realising that it

forms a small part of a large tapestry, the size of which is entirely justified by the overall

340

Hatfield Memorial Lectures Vol. II

technological importance of the subject. I make no apology for dwelling on the more scientific aspects of the formation of ferrite in alloy steels, for, apart from being an area of considerable basic interest, I believe that an understanding of the underlying phenomena is essential to the further development of strong steels that make full use of alloying additions.

ACKNOWLEDGMENTS Much of the work referred to has been carried out in Cambridge by research fellows and students. I would like particularly to thank Drs D. V. Edmonds, P. R. Howell,]. V. Bee, N. K. Balliger, D.]. Walker, P. D. Southwick and R. A. Ricks for discussions, and for their help in providing photographs and data, often prior to publication. I am grateful to the Science Research Council, British Steel Corporation and the Ministry of Defence (Fort Halstead) for support of the research programme.

REFERENCES 1. H. I. AARONSON: York, 1962, 387. 2. K. I. KINSMAN, E. 3. R. K. GOODENOW 4. F. B. PICKERING:

The Decomposition oj Austenite by Diffusional Processes, Interscience,

New

EICHEN and H. I. AARONSON: Metall. Trans., 1975, 6A, 303.

and R. F. HEHEMANN: Trans. AIME, 1965,233,1777. 'Transformation and hardenability in steels', Greenwich, Conn., Molybdenum Company, 109. 5. J. V. BEE and R. W. K. HONEYCOMBE: Metall. Trans., 1978, 9A, 587. 6. 'Micro-alloying '75', Greenwich, CN, 1975, Climax Molybdenum Company.

7. T. N. BAKER: 'Microalloyed steels', Sci. Prog., 1978,65,493. 8. D.]. WALKER and R. W. K. HONEYCOMBE: Met. Sci., 1978, 12,445. 9. A. B. LE BON and L. N. DE SAINT-MARTIN: 'Micro-alloying' 75', Greenwich, Climax Molybdenum Company. 10. P. D. SOUTHWICK and P. R. HOWELL: unpublished

CN,

Climax

1975,

work.

11. P. R. HOWELL and R,' A. RICKS: unpublished work. 12. N. K. BALLIGER and R. W. K. HONEYCOMBE: Met. Sci., 1980, 14, 121. 13. R. A. RICKS, P. R. HOWELL, and R. W. K. HONEYCOMBE: Metall. Trans., 1979, 10A, 1049. 14. K. LOCKE and K. DETERT: Acta Metall., 1957,5,628. 15. M. MANNERKOSKI: Acta Poly tech. Scand., 1964, Ch. 26; Met. Sci.]., 1969,3,54. 16. W. B. MORRISON and]. M. WOODHEAD:]. Iron Steel Inst., 1963,201,43. 17. R. W. K. HONEYCOMBE: Metall. Trans. 1976, 7A, 915. 18. A. T. DAVENPORT and R. W. K. HONEYCOMBE: Proc. R. Soc., 1971, A332, 191. 19. K. CAMPBELL and R. W. K. HONEYCOMBE: Met. Sci., 1974,8,197. 20. P. R. HOWELL,]. V. BEE, and R. W. K. HONEYCOMBE: Metall. Trans. 1979, lOA, 1213.

Ferrite

341

21. 22. 23. 24. 25. 26. 27. 28.

A. RICKS: PhD dissertation, Cambridge, 1979. F. G. BERRY and R. W. K. HONEYCOMBE: Metall. Trans., 1970,1,3279. D. V. EDMONDS:]. Iron Steel Inst., 1972,210,363. N. K. BALLIGER: PhD dissertation, Cambridge, 1977. A. BARBACKI and R. W. K. HONEYCOMBE: Metallography, 1976,9,277. N. K. BALLIGER and R. W. K. HONEYCOMBE: Metall. Trans., 1980, in press. G. L. DUNLOP and R. W. K. HONEYCOMBE: Philos. Mag., 1975,32,61. P. R. HOWELL, P. D. SOUTHWICK and R. W. K. HONEYCOMBE:j. Microscopy, 1979,116,

29.

P. D. SOUTHWICK:

R.

151. PhD

dissertation, Cambridge, 1978.

THIRTY-SIXTH

HATFIELD

MEMORIAL

LECTURE

Clean Steel, Dirty Steel J.

Nutting

At the time the lecture was given Professor Nutting was in the School of Materials} Division of Metallurgy} University of Leeds} UK. The lecture was presented at the University of Sheffield on 6 December 1988.

The embrittling effects of sulphur in relation to the manganese content of steel are discussedin the light of recent research showing the importance of solid state segregation of sulphur to austenite grain boundaries. The influence of sulphur on the overheating characteristics of alloy steels is then described and it is shown that overheating can be prevented by reducing the sulphur and manganese contents. Recent studies on the temper embrittlement of alloy steels have also shown that this problem can be overcome by lowering the manganese content and by reducing the phosphorus and silicon and specifying low antimony, arsenic and tin levels. This research has led to the production of superclean 3.5Ni-Cr-Mo-V steels containing 0.02%Si, 0.01%Mn, O.002%P and O.OOl%Swhich have then been forged to give low pressure steam turbine rotors. The properties of the superclean steels in relation to overheating, temper brittleness, creep, fatigue and stresscorrosion behaviour are described. It it also pointed out that at a given strength level superclean steels have much higher upper shelf energies than conventional steels. The importance of sulphur and phosphorus as deliberate alloy additions to steel is briefly mentioned.

INTRODUCTION In the production of metals we must never lose sight of the fact that they come from the ground. We must also bear in mind that many of the other materials we use in the manufacture of metals also come from the ground. As a consequence the origins of our raw materials greatly influence the final properties of the metals we are trying to extract. The steelmaker has long been aware of these issues and much of his endeavour has been devoted to the problems imposed by the presence of sulphur and phosphorus in his raw materials, and the way these two elements then affect the properties of the steel. Before proceeding further let us just establish a few facts about these two elements in relation to iron. Iron has a crustal abundance of 5.6% while, for sulphur, it is 0.026%, and for phosphorus 0.10%. Thus the percentage crustal abundance ratio ofS/Fe is 0.46 and of P IFe 1.7%. The percentage ratio ofS/Fe and P/Fe in iron ores varies considerably, but is of the order 0.4 and 1.2, respectively. That is to say that during the conversion of primary rock

343

344

Hatfield Memorial Lectures VoL II

into sedimentary iron deposit the sulphur and phosphorus are largely transferred into the ore. There are exceptions depending on the ore genesis and low sulphur and phosphorus content are found. The other raw materials used by the ironmaker also contain sulphur and phosphorus. The early molten iron produced showed little reduction in the content of these two elements and they remained when the iron was converted into steel. Bessemer was well aware of the problems associated with phosphorus and was only able to make his steelmaking process give satisfactory materials ifhe used low phosphorus pig iron. It would seem that he was not fully aware of the difficulties associated with sulphur. But R. F. Mushet.! the son of David Mushet of tool steel fame, certainly was. It was largely through R. F. Mushet showing that the low carbon Bessemer blown steel could be deoxidised by the addition of spiegeleisen, that a sound blow hole free steel could be produced, but as he pointed out, the addition of the spiegel also cured hot shortness. All this happened some 130 years ago, and the consequences are still with us. As every first year undergraduate knows, if sulphur is present in steel it forms FeS which melts at about 1000°C and then forms as a liquid film around the austenite grain boundaries so inducing hot shortness, i.e. lack of ductility during hot rolling. The cure is delightfully simple; add sufficient manganese when MnS is formed in preference to FeS. As MnS melts at 1400°C, no liquid films are formed and so no hot shortness. The history of these developments were reviewed by Melford- in 1980 and will not be recounted further. The question to be answered is how much manganese needs to be added to ensure that the sulphur is 'tied up' as MnS? Research settled the issue and, by the 1920s, it was clearly established that a MnlS ratio of20:1 was required. At that time good basic open hearth steel had a sulphur content of about 0.02% with acid open hearth steel being somewhat higher. To make sure that all steels could be hot worked without cracking, manganese contents of 0.4-0.6% were specified and with remarkably few exceptions such specifications are still with us some 60 years later. But during the 60 years open hearth furnaces have disappeared and sulphur contents as low as 0.001% can readily be achieved by the new steelmaking methods.

EMBRITTLING OF IRON BY SULPHUR Let us move away from the first year undergraduate lectures and see how recent research has provided an answer to the why and how sulphur embrittles iron. The iron-sulphur diagram, shown in Fig. 1, tells us that the sulphur solubility in y-iron at 1000°C is 0.013 wt-% (0.22 at.-%). It could well be asked why the sulphur solution is so low. The atomic diameter of sulphur is 0.212 nm, while iron in the y-form has an atomic diameter of 0.252 nm. Thus the atomic size difference is 15% which, is not too far away according to the Hume-Rothery rules for extensive solid solubility. It is interesting to compare sulphur with our other bad element, phosphorus. The phosphorus atom (diameter 0.216 nm) is slightly larger than the sulphur atom; thus the size difference in relation to iron is about 14%. However, the solubility in phosphorus in

Clean Steel, Dirty Steel

345

liquid 6-FG

y YF~

o Fig. 1

0-01

0-02

0-03 S,wt_o,o

+

F~S + liquid

F~S

0-04

0,05 0-06

Fe-S equilibrium for Fe rich region.

y-iron at 10000e is 1.5 wt-% (2 at.-%). Thus, on an atomic basis phosphorus is about 10 times more soluble than sulphur in austenite. These differences cannot be attributed to the relative stability of the compounds of FeS and Fe3P. Since Fe3P is much more stable the FeS (FeS i\G9 = - 167 kJ mol-l, whereas for Fe3P i\G8 = - 305 kJ mol:"). Let us say the facts are clear, but the explanation of the differences is far from satisfactory. Although it has long been suspected that elements dissolved in a metal can segregate to grain boundaries, it is largely from the work of Hondros and Seah-' that we can quantify the extent of this segregation. In general, the extent of the segregation varies inversely with the solubility of the solute in the solvent metal (Fig. 2). In the case of sulphur in y-iron at 10000e the boundary enrichment ratio is about 104-105 of the amount of sulphur in solid solution. It would now appear that it is the segregated sulphur which causes the embrittlement. This is clearly shown by the work ofTacikowski et al.,4 who found that in pure iron the addition of 89 ppm of sulphur reduced the hot ductility almost to zero at 9800 (see Fig. 3) and this was associated with an intergranular failure mode. This sulphur content is well within the solubility limit at a temperature of 980°C and this temperature is some 8 K below the melting temperature of FeS, as can be seen from Fig. 4. We see, therefore, that one of the 'immutable facts' of ferrous metallurgy is wrong. Hot shortness in the absence of manganese is not a result of liquid films of FeS forming at the grain boundaries of austenite but arises from solute segregation of sulphur in the solid state.

e

ROLE OF MANGANESE There is no doubt that the addition of manganese to steel reduces the susceptibility of hot shortness and it is reasonable to ask how this is brought about.

346

Hatfield Memorial Lectures VoL II

Atomic solid solubility

Fig. 2

Grain boundary enrichment

900

ratio in relation to solid solubility.

1000

TEMPERATURE,

Fig. 3

X Co

3

1'00 ·C

Influence of sulphur on embrittlement of pure iron at high temperature; on each line relate to sulphur contents in ppm."

numbers

Clean Steel, Dirty Steel

347

988 ·C

.u w ~950

4

a:::

w e,

1:

w t-

FS 30 )(

OOOi.

0008

0-012

SULPHUR CONTENT,

Fig. 4

0-016 wt-O/o

0020

Most severe embrittling conditions at different sulphur contents obtained from Fig. 3 shown in realtion to appropriate region of the Fe-S diagram. 4 :

The ternary equilibrium diagram for Fe-Mn-S has not feen fully determined, but the critical part we require ',is known as a result of work carried out over 30 years agC?by Turkdogan et al.5 They calculated the form of the solvus curve for the solubility of sulphur in iron containing differing amounts of manganese. Their results showed that at 1000°C the addition of 0.5% manganese to iron reduces the solubility of sulphur from about 100 ppm to about 1 ppm (see Fig. 5). This comes about because the MnS is more

Stwt-ppm

Fig. 5

Influence of manganese on solubility of sulphur in austenite.>

348

Hatfield Memorial Lectures VoL II

stable than FeS (dG9 MnS = - 370 kJ mol-i). At least this part of our cherished beliefs is intact. It could be argued that decreasing the solubility of sulphur in the solid solution by a factor of 102 should increase the grain boundary enrichment ratio by a similar factor, i.e. manganese should make the embrittling effect worse, but this is a wrong interpretation of the results. In fact manganese, through its affinity for sulphur, desegregates the sulphur and with only 1 ppm of sulphur in solution in the austenite at 1000°C the boundary segregation is insufficient to produce embrittlement but other difficulties arise.

OVERHEATING OF STEEL It has long been recognised that when steels are heated to excessively high temperatures during forging, a form of embrittlement may occur which only becomes apparent when the steel is fully heat treated to put it in its toughest condition. The nature of this embrittlement was established by G. Wesley Austin" in 1936. During the Second World War overheating became a major problem as, in attempts to increase production rates, forging temperatures were raised. Further research was started in an attempt to find the causes of the embrittlement and many papers were published. The whole of this work has recently been reviewed. 7 As the forging temperature of a steel is increased the embrittling reaction gives a decrease in the tensile properties, as shown in Fig. 6, and this is associated with the appearance of intergranular facets on the fracture surface (see Fig. 7). The lowest temperature of preheat to give facets has been defined as the overheating temperature. Inter-

•

, • , •

Fig. 6

Effect of preheating

cooling rata, Kmin-1 2 10 200 2500

temperature followed by full heat treatment fracture for steel EN39.8

on true strain to

Clean Steel, Dirty Steel

349

estingly this temperature decreased as the sulphur content decreased. This was the state of affairs reached in 1948 when, as the wartime difficulties were over, work on the topic stopped. However, by the middle 1960s difficulties with overheating arose yet again, for with the newer steel melting and refining techniques available, sulphur contents were reduced to 0.005% and reheating for forging at 1150-1200°C was sufficient to cause trouble through overheating (see Fig. 8).

Fig. 7

Intergranular

facets on fracture surface of overheated micrograph.

steel; scanning

electron

Fig. 8

Influence of sulphur on overheating temperature of low alloy steels; BE is basic electric; CEV AM is consumable electrode, vacuum arc remelted; 0 H is open hearth. 7

A determined effort was made to track down the causes of overheating and this was so much easier because scanning electron microscopes and their associated analytical facilities were now available. It was found that the embrittlement was caused by the precipitation of MnS on the austenite grain boundaries during cooling after forging. From the Fe-Mn-S

350

Hatfield Memorial Lectures VoL II

diagram it can be seen that at, say, a manganese content of 0.5% the amount of sulphur taken into solution on heating increases to about 0.004% at 1400°C. If the steel is cooled at an appropriate rate the sulphur precipitates as fine particles of MnS on the boundaries and this lowers the impact resistance when the steel is in the toughest condition (see Fig. 9). The ideal cooling rate for precipitation would seem to be within the range 10-200 K mirr", and these are the cooling rates most likely to be encountered during forging (see Fig. 10).

Fig. 9

Sulphide

particles on intergranular electron

fracture surface of an overheated micrograph.

steel; scanning

160

-: (!)

eJ100 Z

w

O~~~~~~~~~~~~~~~ 1

10 COOLING

Fig.vlf)

Influence

102 RATE, Kmin-'

of cooling rate from 1400°C on Izod impact value of fully heat treated 3.5Ni-Cr-Mo-V steel with different Mn contents.f

Clean Steel, Dirty Steel

351

An entirely reasonable question to ask is that if this is the explanation of the embrittling effect, why are low sulphur steels (electric arc) more prone to overheating than the higher sulphur steels (open hearth)? The answer would seem to be that, with open hearth steels because of the high sulphur content, there is always a large number of sulphides still available in the steel after the maximum amount of sulphur has been taken into solution, and these provide sites other than the austenite boundaries on which the dissolved sulphur can be precipitated. A further factor, and possibly the major one, is that if the steel already has a low impact value as a result of containing many manganese sulphide particles within the grains, it needs a more severe weakening effect to occur at the austenite grain boundaries before they become so weak that the ductile fracture path follows around them preferentially.

A CURE FOR OVERHEATING How can we stop a steel from overheating during forging? The Fe-Mn-S diagram is very helpful in providing answers to the question. As can be seen from Fig. 5, we could raise the manganese content, so reducing the amount of sulphur taken into solution and hence being available for precipitation. But this is not an option available to us because increasing the manganese content makes the steel more susceptible to temper brittleness. By the same argument we could add elements to the steel which have an even greater affinity for sulphur than does manganese and, for this, the rare earths (REs) would qualify. But unfortunately the RE additions lead to a dramatic increase in the oxide inclusion content and a general reduction in impact and fatigue properties. Hence this is not a satisfactory route to follow. A feasible, but at first sight outrageous, suggestion would be to reduce the manganese content. Decreasing the manganese content will increase the sulphur solubility and, for a given sulphur content, it will lower the temperature at which sulphur will first precipitate from the austenite and, as can be seen from Fig. 11, the overheating temperature is reduced if the manganese content is lowered for, with a low sulphur and low manganese content, the cooling rate from the forging temperature is not important; the steel just does not overheat (see Fig. 10). If, therefore, we lower both the sulphur and the manganese contents, we may be able to keep all the sulphur in solution and hence have no sulphide inclusions in the steel whatsoever. With O.OOl%S and 0.02%Mn we should be out of trouble, and we are. What about hot shortness? It is really no problem for, as we have seen, no liquid films of FeS will form and the extent of the grain boundary segregation of sulphur is insufficient to cause trouble. However, in the absence of manganese, but in the presence of chromium, chromium sulphides can precipitate at austenite grain boundaries and give rise to an equivalent form of overheating." In order for MnS to precipitate both manganese and sulphur have to diffuse to the austenite grain boundaries. For the precipitation of CrS, chromium also has to diffuse, but it does not diffuse as readily as manganese; hence chromium sulphide overheating is only likely to occur in large forgings. However, this form of overheating will not occur if the sulphur content is kept low.

352

Hatfield Memorial Lectures VoL II 160

,

>..

e

~120

z w ...

U

~

~ 80 o

o ~

Fig. 11

Influence of manganese on Izod impact value of fully heat treated EN39 steel after overheating at 1400°C.8

TEMPER BRITTLENESS The poor impact properties found in alloy steels when they are hardened and then tempered at 500-550°C are so well known that they do not need to be described in detail. As every second year undergraduate knows, having acquired a knowledge of the elegant solution to hot shortness by the addition of manganese, temper brittleness is no problem. All that is required is the addition of 0.5% Mo, a cure developed in the 1920s, but not as efficacious as was originally hoped. Molybdenum is a short term palliative for when steels are exposed to service temperatures of 550°C for prolonged periods the brittleness returns. Modem analytical techniques, particularly Auger spectroscopy, have enabled us to identify the elements in steel which segregate to a-iron grain boundaries and give temper brittleness. These have been found to be phosphorus, arsenic, antimony and tin. The solubility of phosphorus in 'Y-iron has already been discussed but in a-iron it is somewhat higher, reaching a maximum of 3 wt-%. The other elements, arsenic, antimony and tin have solubilities in a-iron within the range 5-10%; that is to say they are readily soluble. It could be asked, in view of the Hondros and Seah-' results, why these elements segregate to grain boundaries and produce embrittlement? In pure iron it is clear these elements do not segregate sufficiently to cause embrittlement but in steels, where we have added manganese to control the sulphur and to help with deoxidation and made further additions of silicon also to help with deoxidation, there is an interactive effect. The manganese and silicon force the phosphorus, arsenic, antimony and tin to the grain boundaries and embrittlement occurs. After painstaking study Watanabe and Murakami-? showed that the susceptibility to temper embrittlement could be related empirically to the] factor where:

] =

(Mn

+ Si)(P + Sn)

Clean Steel, Dirty Steel

353

A value of] equal to 10-2 gives rise to temper embrittlement in 2.25Cr-1Mo steels, whereas a value of 10-3 would embrittle a Ni-Cr-Mo-V steel. Further work has shown that the formula can be extended to include arsenic and antimony together with the other elements and give them all some weighting. The K value proposed by its originators, Kohno et al., 11 is given by: K

=

(Mn + Si)(10P + 5Sb + 4Sn + As)

From this equation it can be seen that if we wish to avoid temper brittleness we must try

to keep phosphorus levels low and the lesswe have of that element, which we all thought to be beneficial to the steel manganese content, the better.

CLEAN STEEL In giving a critical assessmentof clean steel, Kiessling'< in 1980 made the statement that 'The cleanness of steel, like beauty, is very much in the eye of the beholder', and then added 'The clean steel of yesterday is not the clean steel of today'. At that time cleanness was seen as the absence of sulphur, phosphorus and oxygen together with the trace elements arsenic and antimony coming from the original raw materials and the tramp elements copper and tin picked up from scrap. However, from the evidence which has been presented we must now include among our list of proscribed elements manganese and silicon, particularly if we are to make low alloy steels which will forge without overheating and which will be free from temper brittleness during initial heat treatment or after prolonged exposure to temperature in the range 450-550°C. The Electric Power Research Institute (EPRI) of the USA (Re£ 13) had confidence in this assessment to the extent of investing some $8m in a programme of research and leading to the production of superclean steel forgings for fabrication into low pressure steam turbine rotors. The basic composition used has been that of a 3.5Ni-Cr-Mo-V material. The contents of the proscribed elements in these forgings is as follows: Mn Si S P As Sb Sn

wt-% 0.02-0.05 0.02-0.03 0.001-0.0015 0.002-0.003 0.003 0.001 0.003 ] value

=

5-8

X

10-4

354

Hatfield Memorial Lectures VoL II MAKING A SUPERCLEAN STEEL

The steels have been made initially in the basic electric arc furnace with an oxidising slag to remove phosphorus, manganese, and silicon. This is followed by ladle refining with a reducing slag to remove sulphur. The oxygen is then removed by carbon in a vacuum furnace. Hydrogen and nitrogen have been removed by argon bubbling and by vacuum stream degassing while casting the ingot. Seven 90 t casts of superclean steels have been produced. Two have been made in Austria and have then been forged into low pressure rotors before being sectioned for testing. The remaining casts have been produced in Japan and four of these have been forged and heat treated to give low pressure rotors which are to be installed in power stations in the USA and Japan. A finished rotor is shown in Fig. 12. To produce such a superclean rotor an overall premium of 20-30% has been suggested above the cost of a conventional rotor steel forging.

Fig. 12

Low pressure rotor produced

PROPERTIES

from superclean steel forging.

OF SUPERCLEAN STEEL

Superclean steels, when examined in the optical microscope, show virtually no nonmetallic inclusions. The area fraction is about 0.004-0.008% going from core to rim. But having a classical MnlS ratio of20:1 any sulphides that would form are MnS. On ductile fracture surfaces the scanning electron microscope reveals only very tiny MnS particles (Fig. 13). However, the extent ofMnS precipitation which can occur at these manganese and sulphur levels is insufficient to cause overheating, as can be seen from Fig. 14. In fact the Charpy impact value in the fully heat treated condition (quenched from 840°C and tempered at 600°C) increases as the preheating temperature is increased. All the sulphides are taken into solution and then, on cooling, the cooling rate through the critical temperature range where sulphides could precipitate is so rapid in relation to the temperature that no precipitation occurs. Superclean steels do not show temper brittleness. This is clearly shown in Fig. 15, where it can be seen that aging for periods of up to 10 000 h at temperatures between 350

Clean Steel, Dirty Steel

355

Fig. 13 Transgranular fracture surface of superclean 3.5Ni-Cr-Mo-V steel in fully heat treated condition but after pretreatment at 1400°C and air cooling, no overheating; scanning electron micrograph.

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356

Hatfield Memorial Lectures VoL II 100~--------------------------------~ Temp.,"e

~ j~g v

o

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c

400

425

450 475

centre

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and 475°C shows very little change in the fracture appearance transition temperature. The higher transition temperature for samples taken from the centre of the ingot than for those taken from the surface is what would be expected from the difference in microstructures which occur as a result of differing cooling rates at the surface and centre of the ingot. There are other improvements in the properties of superclean steels over those of conventional purity steels which arise from the lower inclusion content of the superclean materials. Creep failure in rotor steels is triggered off by the nucleation of voids which develop from sulphides at the austenite grain boundaries. If there are no sulphides at the grain boundaries void nucleation will still occur but it will take longer at a given temperature or, in a given time, will occur at higher stress levels. The general improvement in creep properties is shown in Fig. 16. Fatigue failure is frequently nucleated by inclusions and therefore superclean steels have improved fatigue properties. This argument applies particularly to low cycle fatigue where the improved ductility of the low inclusion content steel gives rise to improved low cycle fatigue properties, as can be seen from Fig. 17. The stress corrosion properties of superclean steels are also improved insofar as the absence of impurities leads to a significant increase in the time for crack initiation at a given stress and for about a 10% increase in the stress for crack initiation within a given time (see Fig. 18). However, the crack propagation rate is not increased when the impurities are removed. The role of inclusions and carbides in controlling the ductility of steels is well known from some of the pioneering work of Gladman et al.14 (see Fig. 19). Since that time Briant

Clean Steel, Dirty Steel

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It might be expected, therefore, that reducing the sulphur and the phosphorus contents, and also lowering the contents of the elements such as manganese and silicon which have a synergistic effect on phosphorus segregation, would lead to a considerable improvement in the impact properties of steels. This is clearly seen to be the case as a result of work carried out by Kearney. 16 The impact values of a superclean steel when tempered to give high strength and hardness levels is about twice that found in the same steel of conventional composition (see Fig. 22). It could be argued that while a Charpy test is of interest, the real question to be asked is does the removal of the proscribed elements improve the fracture toughness? The answer

Clean Steel, Dirty Steel

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.

(a) 850°C, 1 h; (b) 1200°C, 3 h.

360

Hatfield Memorial Lectures VoL II

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would certainly seem to be 'yes' and as can be seen from Fig. 23, for a given yield strength the fracture toughness is significantly improved.

FREE CUTTING STEELS, PIANO WIRE AND NEO-CONFUCIANISM Having thoroughly damned sulphur in relation to its influence on properties such as toughness, there are times when we do not need tough steels, particularly if we want good machinability. I am sure it will come as no surprise to say that Henry Bessemer discovered that the addition of sulphur to steel produced an improvement in machinability. By the beginning of the 20th century the production of high sulphur, free cutting steels was a well established business and the steel stockists's catalogues gave specifications for free cutting steels in a range of rolled bar sizes. Much has been written about how the machinability is improved by the addition of large volume fractions of MnS to steel. Recent research on the high strain deformation of

Clean Steel, Dirty Steel

361

steels is beginning to shed further light on the structural changes which occur in machine chips and hence on the machining process.17,18 The other element which has been our major concern, phosphorus, was not always looked on as a dangerous element to have in steel. Recent historical research has shown that in the 18th century the wires used for stringed percussion instruments were low carbon steels containing up to 1% phosphorus.tv-e? It would appear that phosphorus dissolved in ferrite increases the work hardening modulus and, at the same time, hinders high stress relaxation; a vital requirement if an instrument is to stay in tune. It was only in the middle of the 19th century that these wires were replaced by patented medium carbon heavily drawn wires; i.e. modern piano wires. Deliberate phosphorus additions are now being made to certain steels for producing rolled sheets. It would seem that the elements which find their way into steels are like the rest of us; we have both good and bad characteristics! In the terms of the later Chinese philosophers, the Neo-Confucianists, these elements have their yin and their yang. When they are yin they are impurities and when they are yang they are deliberate alloy additions. The yin qualities of sulphur and phosphorus are well recognised, but it has taken some 130 years to appreciate that after Bessemer and Mushet extolled the yang qualities of manganese, this element could also be yin. We now know that if we want to make strong, tough, heat treated, low alloy steels, we must keep out sulphur, phosphorus, manganese and silicon.

ACKNOWLEDGEMrnNTS I was first introduced to the intricacies of alloy steels by the later Professor A. Preece when, in 1944 as an undergraduate, I started work on a project on the overheating and burning of steel. I am ever grateful for his introduction and his later encouragement. Since that time many organisations have provided funds for me to continue work on the topic. I would particularly like to thank Dr R. I. Jaffee of the Electric Power Research Institute (EPRI), Palo Alto, CA, USA, for financial support during the last eight years and for an opportunity to take a small part in the EPRI Clean Steel Project.

REFERENCES 1. R. F. MUSHET: 'The Bessemer Mushet process or manufacture of cheap steel'; De Archaeologische Pevs, Nederland. (Reprinted by The Historical Metallurgy Society, 1984.) 2. D. A. MELFORD: Philos. Trans. R. Soc., 1980, A295, 89-103. 3. E. HONDROS and M. P. SEAH: Int. Met. Rev., 1977,22,262-301. 4. M. TACIKOWSKI, G. A. OSKINOLU and A. KOBYLANSKI: Mater. Sci. Technol., 1986,2,154-

158. 5. E. T.

TURKDOGAN,

S. IGNATOWICZ and].

PEARSON:].

Iron Steel Inst., 1955, 180,349-354.

362

Hatfield Memorial Lectures Vol. II

6. G. W. AUSTIN: in 'First report of the Alloy Steel Committee', The Iron and Steel Institute, London, 1936, 189-211. 7. G. E. HALE and]. ,NUTTING: Int. Met. Rev., 1984,29,273-298. 8. N. P. MACLEOD and]. NUTTING: Met. Technol., 1982,9,399-404. 9. B.]. SCHULTZ and C.]. McMAHON,]R: Metall. Trans., 1973,4,2485-2489. 10. ]. WATANABEand Y. MURAKAMI:Am. Petrol. Inst., Preprint no. 28-81,1981,216-224. 11. H. KOHNO, M. MIYAKAWA,S. KINOSHITA and A. SUZUKI in Advances in Material Technologiesfor Fossil Power Plants, ed. R. Viswanathan and R. 1.] affes, ASM International, Metals Park, OH, 1987,81-88. 12. R. KIESSLING:Met. Sci., 1980, 14,161-172. 13. Electric Power Research Institute: in Proc. Meeting 'Super clean 3.5%NiCrMoV and improved CrMo V steels', Dusseldorf, June 1988, Verein Deutscher Eisenhiittenleute. 14. T. GLADMAN, B. HOLMES and 1. D. McIVOR: in Effect of Second-Phase Particles 011 the Mechanical Properties of Steel, 1971, The Iron and Steel Institute, London, 1971, 68-78. 15. C. L. BRIANT and S. K. BANERJI:Metall. Trans., 1981, 12A, 309-319. 16. M. G. KEARNEY:Unpublished work, University of Leeds, 1988. 17. R. L. AGHAN and]. NUTTING: Met. Technol., 1981,8,41-45. 18. M. BINGLEY:PhD thesis, University of Leeds, 1984. 19. M. GOODWAY: Science, 1987,236,927-932. 20. M. GOODWAY and R. M. FISHER:]. Rist. Metall. Soc., 1988,22,21-23.

Author Index Bain, E. c., 45-94 Bastien, P. G., 95-116 Darken, L. S., 247-282 Honeycombe, R. W. K., 315-342 Hume-Rothery, W., 193-222 Kurdjumov, G. V., 117-162 MeW, R. F., 7-44 Menter,]. W., 283-314 Morrogh, H., 223-246 Nutting, ]., 343-362 Quarrell, A. G., 3-4,163-192

363

Subject Index a-solid solution, 146-152

duplex stainless steel, 337-339

adsorption, 258-263 alloy systems, 195,213-216 atomic diameter, 200 Auger spectroscopy, 293 austenite, 7-44, 321-325 austenite/ferrite interface, 337-339 austenite-ferrite transformation, 316-320 austenite-martensite transformation, 129-145

elastic stress, 268-272 electromagnetic lenses, 182 electron metallography, 60-65 electron probe microanalysis, 177-180, 289-290 embrittling by sulphur, 344-345 eutectic carbides, 231-233, 236-238 eutectoid steels, 11-19, 20-23 fatigue behaviour, 73-74 ferrite, 23-31, 112-113, 315-342 grain size, 320-323 fibrous carbides, 331-333 field ion microscopy, 248-251, 290-292

bainite, 7, 35-41, 319 banded structures, 95-116 appearance, 96-102 mechanism of formation, 106-113 boron effect, 59 British Iron and Steel Research Council, 3-4 brittle fracture, 75-82

grain boundary segregation, 289-293 graphite, 227-231, 238-244 growth processes, 7-44 hardness of quenched steel, 125-128 Hatfield, William Herbert, 3-4, 45-46, 95, 193 heterogeneity of cast steel, 294-312 heterogeneity of steel, 283-314 hydrogen interactions, 277-279 hypo-eutectoid steels, 8, 23-33

carbide phase, 152-154 carbide precipitation, 326-335 cast irons, 223-246 mechanical properties, 238-243 solidification, 224-238 clean steel, 343-362 constitutional segregation, 294-299 crystal structure, 205-212, 216 crystals, 68-70 filamentary, 70-73 spiral growth, 68-70

inclusion origin, 300-302 intennetallic chemistry, 193-222 interphase precipitation, 329-331 interstitial solutions, 254-258, 268-272 isothermal reaction, 8-10

dark field illumination, 187 deformation, 75-82 dendritic segregation, 102-106, 111-113 diffilsion mechanism, f>5-6f>

liquid solutions, 251-254 low carbon steel, 119-122

dirty steel, 343-362 dislocations, 66-68 dispersion stability, 335-337

machinability, 74-75 manganese, 345-348, 350

365

366

Hatfield Memorial Lectures VoL II

martensite, 53-56,118-119,122-125 metallic chemistry, 247-282 microalloyed steels, 339-340 microradiography, 103 nitrogen interactions, 272-277 non-destructive testing, 309-312 non-metallic inclusions, 299-312 nucleation, 7-44 overheating of steel, 348-351 pearlite, 11-23,31-32, 165, 181 periodic table, 193-194, 198-199, 205, 216-217 phase changes, 193-198 physical chemistry, 86-89 pro-eutectoid reactions, 7, 8, 15,23-25, 33-35 quantitative autoradiography, 103 quantitative metallography, 171-172, 175-176 quenching,117-162 reaction rates, 264-268

residual elements, 285-289 Scheil analysis, 173, 175 solid solutions, 251-254 solid-liquid transformation, 199-204 Sorby, Henry Clifton, 3-4, 163-192 superclean steel, 354-360. supersaturated ferrite, 333-335 surface hot shortness, 285-289 temper brittleness, 352-353, 357 tempering, 56-59,117-162 thermodynamics, 82-86 thin film microscopy, 184-186 transformation kinetics, 323-326 transmission electron microscope, 180, 184-188 USA, trends in research, 45-94 Widmanstatten pattern, 24-28, 33, 165, 168-169,176,318,338 X-ray diffraction, 170 X-ray fluorescence, 52-53, 103-104 X-ray metallography, 48-52