The Hatfield Memorial Lectures Volume II
Dr William Herbert Hatfield FRS, 1882-1943. An informal picture taken in 1942...
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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.-
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a.
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~
V
Vi
J
4 5 b 2 I 3 CARBON CQ.JCENTRATION ATOMIC
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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.
-,
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"
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
153·0
154·0 155·0
DI FFRACTION
ANGLE
technique
15b·0
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157·0
degrees of arc
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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Fig. 8
X-ray fluorescence analysisapparatus.
crystal spectrometer for simultaneous analysis of chromium, steel.
nickel and molybdenum
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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co!.PresSion r- 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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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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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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Strippinq Section
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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.
--..-
Tarqet for condensinq beam
Cell
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
/
IL 1/
0'8 ,I)
/
/
/
/
./
/ ~/
'" O·b
U
1/
>-
::
?
/ .(Raoult's
>
law
1/
;:: 0-4
u 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
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)
~
-l
~
0
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+200
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8
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TEMPERATURE,OC
-lao
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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
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Phenomena Occurring in the Quenching and Tempering of Steels
151
50
I
u
O·IO%C. 0.50% Ti
a::
en 40 V')
o·/l%C
w Z
o a::
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~
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20
Fig. 28
Influence
'\
-0-
\
/~
l
\
200 400 600 ANNEALING TEMPERATURE.oC
of Ti on the decomposition of martensite O.1%C.1S
on tempering
steel with
4·0 E E
i
~ ~~
3'2 r-----=---f---+---+---..:..o.:r-----1'r-
« UJ
ex: co
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Tempering
200
ANNEALING
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
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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
166
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)
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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
176
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.
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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
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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
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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
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
'< 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
~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~_
.
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
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' 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.
~ ZZ -~ z~~ 0< ~u
zg
+
0-15
0-10
0-05
V> (5
0
-- ---
------------
0-01 0-02 0-03 NITROGEN IN ANNEALED an, wt - 0/0
--- -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
Ib Iin2
3000 4(X)() 5000
E50
i'::/~ 0. 0.
----------
~ 20 , I
10
o
15
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
40.- __ ~2_0_0_0~0 __ ~ __4_0_00~0 __ ~_6_0_0~0_0 __ ~
1000
Fig. 45
2000
3000
PRESSURE OF H2,atm
4000
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~~
/ /0
/' ,/
/
-
~
0/ /
/
-: o
Fig. 46
Influence
/1
/0 1 0'2
of hydrogen
I
I 0'4 HYDROGEN
I
I I 0'6 PRESSURE,atm
on crack propagation, Williams and N elson>").
I 0-8
I
1'0
low alloy steel, 23°C (data of
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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Hatfield Memorial Lectures VoL II
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
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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