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HIGH SUBCRITICAL TEMPERATURES
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5178
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HIGH SUBCRITICAL TEMPERATURES
P R O E F S C H R I F T
TER VERKRIJGING VAN DE GRAAD VAN DOCTOR IN DE TECHNISCHE WETENSCHAPPEN AAN DE TECHNISCHE HOGESCHOOL DELFT, OP GEZAG VAN DE RECTOR MAG-NIFICUS IR, H.R. VAN NAUTA LEMKE, HOOGLERAAR IN DE AFDELING DER ELEKTROTECHNIEK, VOOR EEN COMMISSIE AANGEWEZEN DOOR HET COLLEGE VAN DE-KANEN TE VERDEDIGEN OP WOENSDAG 4 OKTOBER 1972 DES NAMIDDAGS TE 2 UUR
DOOR
George Benjamin B a r k a y
MECHANICAL ENGINEER, DIPLOM TECHNICAL UNIVERSITY BUDAPEST
GEBOREN TE NYIREGYHAZA (HONGARIJE)
iSRAELISCH STAATSBURGER SEDERT 1963
I0 6X ^JJ^
Printed by M. Weinberg, Haifa,
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BIBLIOTHEEK TU Delft P 1902 5178
TO EDITH AND GABY,
Development Foundation and to Ing. S. Golan, Director of the Israel In-stitute of Metals, for having graciously accorded me the opportunity to carry out this work,
My thanks are also due to Mr. Gideon Schmuckler for his help in overcoming the language difficulties and in presenting this work in its present form.
Finally, I am most grateful to the staffs of the Department of Metallurgy at the Technological University of Delft and of the Foundry Laboratory, Israel Institute of Metals, Haifa, for their help in the exe-cution of the experimental part of this work.
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1. I n t r o d u c t i o n 7
1.1 T h e p r a c t i c a l and t h e o r e t i c a l backgrounds of the
i n v e s t i g a t i o n 7 1.2 Purpose and scope of the research 8
2. D i s c u s s i o n o f p r e v i o u s s t u d i e s 10 2.1 General 10 2.2 G r a p h i t i z a t i o n 11 2.2.1 N u c l e a t i o n and g r a p h i t i z o t i o n in c a s t iron at high temperatures 11 2.2.2 G r a p h i t i z a t i o n of steel at s u b c r i t i c a l temperatures 14
2.3 Growth c h a r a c t e r i s t i c s of cast iron 16 2-4 P r e v i o u s own work on the subject 23 3. E x p e r i m e n t a l w o r k a n d m e a s u r e m e n t s 26
3.1 K i n d s of cast iron used in the experiments 26
3.2 D i l o t o m e t r i c measurements 27
3.3 D i f f u s i o n couples 31 3.4 D e n s i t y measurements 37 3.5 Hardness changes 40 3.6 M i c r o s t r u c t u r a l changes 4 2 3.7 Mossbauer effect measurements 45
3.7.1 D e s c r i p t i o n of the spectrometer 45 3.7.2 Preparation of specimens for Mossbauer measurements 47
3.7.3 Interpretation of the Mossbauer spectra 47 4. D i s c u s s i o n o f t h e g r o w t h k i n e t i c s 56 5. D i s c u s s i o n o f t h e d i f f u s i o n c o u p l e s t u d y 63
6. D i s c u s s i o n of the i n f l u e n c e of graphite on dimensional growth 67
6.1 C a l c u l a t i o n of the t h e o r e t i c a l graphite volume 67 6.2 C o r r e l a t i o n and "adjustment of c a l c u l a t e d and measured values 70
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7 . D i s c u s s i o n o f t h e d i f f u s i o n p r o c e s s 77
7.1 Carbon diffusion 77 7.2 Vacancy diffusion 78 7.3 The influence of silicon 80
7.3.1 The effect of silicon on carbon diffusion 81 7.3.2 The effect of silicon on vacancy diffusion 84 7.3.3 The question of the dissolution of silicon in cementite 85
8 . G e n e r a l d i s c u s s i o n 87 8.1 Dimensional growth theory at subcritical temperatures 87
5.1.1 Pearlite decomposition 88 8.1.2 Excess growth due to the shape of graphite particles 92
8.2 Further theoretical considerations of the diffusion process 93
8.2.1 General concepts of the diffusion 95 8.2.2 Application of the general concept to the present cose 99
8.3 Final interpretation of the results 107
9 . S u m m a r y 111 l O . S a m e n v a t t i n g 114
1 1 . R e f e r e n c e s 116
1. I N T R O D U C T I O N
1.1 T H E P R A C T I C A L AND T H E O R E T I C A L BACKGROUNDS OF THE I N V E S T I G A T I O N
It is a fact well known to investigators and users of cast irons that account has to be taken of dimensional growth at elevated temperatures. This phenome-non is of considerable importance, since cast iron is widely used at such temper-atures. From the point of view of dimensional growth at isothermal heating these temperatures may be subdivided into three ranges :
(a) Below about 400°C. where there is practically no growth;
(b) Above the critical temperature, where the dimensional growth occurs mainly as a result either of the transformation of pearlite to austenite and graphite or of the ferrite-austenite phase changes;
(c) Above 400°C but below the critical temperature.
FYom the above it is clear that, in practice, service at temperatures as de-fined under (a) does not pose any problems, while temperatures above the critical point — see (b) — are not as a rule encountered in general practice, their import-ance lying principally in their constituting the reinge in which heat treatments are administwed. It is, therefore, range (c) that is the most problematical as far as the use of cast iron at elevated temperatures is concerned. The critical point re-ferred to is the A, temperature, i.e. the temperature of transition between pear-lite and austenite.
Theoretically, the Aj temperature is 738°C*'\ according to the Fe-C equilibrium diagram. In commercial cast irons this point may vary under the in-fluence of silicon, which raises the critical point, or of other elements, some of which raise it while others lower it.
Obviously, in range (c) the rate change — that growth rate — is the more rapid the higher the temperature.
The principal impetus to the present investigation was given by a previous work of the author of this thesis*^\ in which he examined the factors influencing the service life of ingot moulds. Measurements of the mould temperature showed
that with the exception of the working face, the wall temperature nowhere exceeds 700°C. Thus the growth and the phase changes occuring at temperatures up to that point were taken to exercise the predominant influence on the service life of the moulds, and this was borne out by the experiments. Besides in ingot moulds, cast iron is used in the 400—700°C temperature range in engine blocks, and in some other applications where it undergoes cyclic heating or is working in iso-thermal conditions. The temperature of 700°C was accordingly selected for the present study of the behaviour of cast iron at elevated temperatures, because it is a relatively elevated temperature yet still controllably subcritical. The processes taking place can be subjected to tests of relatively short duration and the results are of value for practical purposes, as they undoubtedly are also applicable to lower temperatures, with the sole exception that they take longer.
1.2 PURPOSE AND SCOPE OF THE RESEARCH
Cast iron is one of the most ancient kinds of iron-carbon alloy and is still the most frequently used of all cast metals. This is why it has retained the in-terest of foundrymen to this day and has remained a fit subject for research by metallurgists side by side with the most sophisticated alloys.
One of the less well known features of cast iron metallurgy is its behaviour at high subcritical temperatures. It was the aim of the present research to sup-plement the practical observations and to provide theoretical explanations for the various phenomena, thereby facilitating the selection of the type of cast iron most suitable for the particular application envisaged.
It was found that the most basic characteristic of cast iron behaviour at subcritical temperatures is the dimensional growth of the iron. It affects the service life of the casting and permits transformation kinetics to be studied quantitatively. Complementary experiments were carried out in order qualitative-ly to evaluate the processes taking place and to propose explanations for them. One of the principal final conclusions to be drawn was that the behaviour (or the service life) of cast iron at high subcritical temperatures can be shown to be de-pendent on two variables, viz. on the phase-change rate and on the overall growth.
The former is a function of the composition of the cast iron or, more precisely, the composition of the cementite and ferrite phases, and is controlled in the first stage by the graphitization and during the second stage by the decomposition of the cementite. The overall growth, on the other hand, is a function of the com-bined carbon content and of the graphite shape or, in other words, what may be^ termed the "form factor" of the graphite. This will be true on condition that the ambient atmosphere exerts but a negligible influence (oxidation, etc.) or is pre-vented from doing so.
The research was mainly concerned with flake graphite cast iron of differ-ent compositions, and only in certain cases, when it was judged advisable, were comparisons made with spheroidal cast irons.
2. D I S C U S S I O N O F P R E V I O U S S T U D I E S 2.1 G E N E R A L
Scientists investigating the behaviour of cast iron at high temperatures have focused their attention especially on the phenomena •
— Graphitization — Dimensional growth
An understanding of the problems involved is of considerable importance. but it is hampered by numerous theoretical difficulties. As to the practical as-pects. graphitization is a phenomenon encountered mainly in the production of malleable cast iron and. with certain reservations, in the manufacture of graphit ized steel, while growth causes trouble principally at high working temperatures.
— G r a p h i t i z a t i o n has been investigated by several authors from two principal points of view ;
(a) nucleation and graphitization in malleable cast iron at or above the critical temperature.
(b) graphitization in steel at subcritical temperatures.
— T h e d i m e n s i o n a l g r o w t h characteristics of cast iron and the factors influencing the mechanism of growth have been studied in the following conditions;
(a) at subcritical temperatures;
(bi in the austenitic temperature range.
(c) at the above mentioned temperatures, in conditions of isothermal heating;
(d) during cyclic heating from subcritical through the critical to high temperatures.
2.2 G R A P H I T I Z A T I O N
2.2,1 Nucleation and Graphitization in Cast Iron at High Temperatures. The two graphitization reactions'^^ to be considered here are ; (a) first-stage graphitization;
Cementite -» Austenite -^ Graphite (b) direct second-stage graphitization;
Austenite -> Ferrite - Graphite
The relevant investigations, designed to observe graphitization and nucle-ation, were carried out at temperatures above the critical point.
To obtain first-stage graphitization the specimens were heated to up to about 900°c'^\ while direct second-stage graphitization was observed after cool-ing from about 900*^0 to a high sub-critical temperature (about 650°C), at which the specimens were quenched.
From metallographic observations some investigators'*'^'*' have conclud-ed that in first-stage graphitization the temper carbon nucleates preferentially either at the interfaces between austenite and eutectic cementite or at the sur-face of inclusions. This interfacial nucleation occurs because the energy as-sociated with an interface lowers the energy barrier to the formation of a critical nucleus.
Furthermore, the greatest concentration of carbon atoms obviously exists at grain boundaries and other internal interfaces, where there is thus a greater probability (than in the body of either a cementite or an austenite grain) of find-ing clusters of sufficiently large numbers of tightly packed carbon atoms for forming graphite nuclei.
Other authors*^'^' have suggested that the nucleation sites are located at the interfaces between austenite and small carbide particles left over in the austenite from the incomplete dissolution of pearlitic cementite. In that theory the small carbide particles are envisaged as growing to some equilibrium size
by drawing carbon from the larger particles of eutectic carbide until graphite nucleation occurs.
Yet other authors have considered, however, that growth of small particles at the expense of larger ones is completely contrary to the theoretical and ex-perimental knowledge of the coarsening of precipitates. They found''' that, with the possible exception of the situation at very low graphitization temperatures, the nuclei are always associated with massive eutectic cementite. It was found'*' that iron having a low carbon content would contain large amounts of Widmanstat-ten pro-eutectoid cementite within the ausWidmanstat-tenite. At relatively low graphitization temperatures, say 800°C, most of that cementite would at first remain out of solu-tion and only gradually spheroidize during the incubasolu-tion period. Structurally there would be no differences between the surface of the newly formed spheroidal particles and that of the eutectic cementite, and nucleation is equally probable on both. At higher temperatures, of course, the pro-eutectic carbide would be in solution and nucleation accordingly confined to the eutectic cementite.
Several authors have studied the effect of alloying elements on the decom-position of cementite and on first and second-stage graphitization reactions in cast iron. A wide-ranging investigation was made by Sandoz'' on the influence of Cu, Ni, Cr, V, and of the combination of Mn and S. The effects of these and other elements were also subjects of published studies'*'^"'^^'^^''^'.
c u ; A number of investigations have shown that Cu accelerates first-stage graphitization'^'^'". Several authors'^' recognized a sigmoid shape in the curves on plots of per cent graphitization against the log of time and found that increasing the amount of Cu shifts these curves to the left, i.e. along the log time axis.
The work of Sandoz'^' shows that increasing the Cu content from 0.08% to 1.21% shifts the sigmoidal curves to the left linearly with the amount of Cu content. The relevant isothermal treatment was given at a temperature of 899°C (1650°F), Sandoz held that the in-fluence of Ni is similar to that of Cu
Mn&
S; The combination of Mn and S, both graphitization-retarding ele-ments. was found to retard both first and direct second-stage graphit-ization. On this point the authors cited agree. The concentration of Mn in the cementite of irons with excess Mn increases slightly dur-ing the reaction. In white cast-irons, a high Mn content retards the graphtization of isolated cementite particles.
Cr; It has been recognized by several authors that Cr strongly retards both first-'^^••^^^^' and second-stage graphitization*^^•^^'. It was prov-ed by Sandoz*^' that the sigmoid curves characterizing graphitization shift to the right as the Cr content i s increased. Investigations cited by Sandoz proved that in white cast-iron the Cr content of the cementite increases during the first-stage graphitization reaction. Further observations by Sandoz' ^* show that a sufficient amount of Cr would cause the cementite to become thermodynamically stable with respect to graphite, and thus impossible to graphtize. Generally speaking, the effects of V are similar to those of Cr.
Sandoz'^' has derived an equation based on the assumption that the solution rate or stability of cementite controls the rate of first-stage graphitization. San-doz's equation is similar to that developed in another form by Schwartz'^', and it r e a d s ;
3 3 2 2
Gf - 8A (t - tg) - 12A (t - tg) + 6A (t - t^)
Where G, = graphitization fraction ^% graph)
' 100 A = constant related to solution rate of cementite
t = first-stage graphitization time tg - incubation time
If it is assumed that t = 0 , then different values of A produce a family of signoid curves which are quite similar to many of the experimental curves shown in Sandoz's investigation.
It is noteworthy that Sandoz himself has observed that the mathematical expression for first-stage graphitization is insufficient for fully specifying the effects of the alloying elements which retard first-stage graphitization. This is due to the fact that increasing concentration of these elements in the cementite during the reaction would progressively lower the value of the solution rate parameter " A " , which is thus seen to be a variable quantity.
Most of the investigators have mathematically characterized the first-stage graphitization reaction by the sigmoidal rate-equation^** '':
Y = 1 - e x p (-Kt") vrtiere Y = fraction transformed in time " t " .
n - constant referred to as time exponent index. K = temperature-dependent rate constant.
In first-stage graphitization the values of " n " were found to be in the range of 1.5 to 4.o'*'^'', while " K " depends on temperature and composition and lies between 8 x 10~* and 13.8~^ '^''.
2.2.2 Graphitization of Steel at Subcritical Temperatures.
Several investigations have been concerned with the graphitization reaction of steel. Generally speaking, it was found that the optimum graphitizing tempera-ture is eSO^C*^^'. The presence of a critical amount of Al and/ or Si and a minimum of carbide stabilizer elements such as Cr, Ti, V, etc. has been shown to be of great importance. It was also found that the graphitization rate of steel is higher after quenching than from the pearlitic structure without quench-ing.
Graphitization is specially accelerated by the alumina dispersions formed by internal oxidation during graphitization'^'. In steel. Si is l e s s active an ac-celerator than it is in cast iron, it has also been suggested that the graphitiza-tion rate is affected not only by the quantity of the alloying elements, but also by their distribution.
Dennis'^' found that hypo-eutectoid steel graphitizes at 600°C in air, whereas in vacuo or in a nitrogen atmosphere at the same temperature only small amounts of graphite are obtained. He suggests that the accelerated graphitiza-tion in low-carbon steels deoxidized with excessive amounts of aluminium results from the heterogeneous nucleation of graphite caused by dispersions of alumina formed m the solid state.
Dennis further holds that the "secondary aispersion". formed by internal oxidation of steel during graphitization is about 30 times more effective than the "primary dispersion" formed during the cooling of the steel. He found that the graphite nodules formed exclusivelj in the ferrite grain boundanes, where diffusion of oxygen should most readily occur.
Similar conclusions were reached by Hickley and Quarrell'^', who investi-gated steel graphitization at 660°C. They found that graphitization is much slower when the treatment is given in vacuo than in the presence of oxygen and that atable steels are particularly sensitive to the partial pressure of oxygen in the atmosphere prevailing during treatment. In unstable steels, this effect is smaller or even negligible. Quenching to martensite before giving the treatment at 660°C accelerates graphitization. and so. but less effectively, does previous cold-working. Hickley and Quarrell established that presence of metallic alumi-nium or of silicon in major quantities promotes graphitization. They suggest that these phenomena can be explained by the production of y-alumina or cris-tobalite When steels containing aluminium or silicon are treated at 660°C in the presence of oxygen, their oxides act as nuclei for graphite.
The kinetics of steel graphitization was investigated by Rosen and Dirn-feld*^'^^' in different heat treatment atmospheres and at different temperatures. They found that the sigmoidal rate equation - used for describing the first-stage graphitization reaction of cast-iron — is also suitable for steel graphitization. The value of the time exponent " n " was found to be between 2 and 2.5, and " K " - the temperature rate constant - between 0.55 x 10~^ and 6.59 x 10~^ at lower temperatures'^* or between 0 151 and 1.245 x lo~* '^' at higher tempe-ratures ranging from 575"'C to 700°C . The difference in the K-values obtained
is due to the differences in pretreatment given by the two authors. In special conditions of previous quenching and tempering " n " = 1.90 - 2.98 and " K " =0.0193 X 10-^ to 83.60 X 1 0 - ^
2.3 G R O W T H C H A R A C T E R I S T I C S O F C A S T I R O N
The structural changes in cast-iron during heating are made manifest by what is known as growth. Much work has been invested in measuring the charac-teristics of growth and to find an explanation for it. For these purposes, tests were made both below and above the critical temperature, with both isothermal and cyclic heating and in different atmospheres.
Several of the observations concerning structural changes have been made by measuring hardness as a function of time, various types of cast-iron being compared.
Other authors used dilatation measurements to determine growth charac-teristics.
Gilbert and White'^*' tested the behaviour of flake and nodular cast=irons with and without added tin during isothermal heating in the temperature range 550°—650°C. The original, as-cast structure was more or less pearlitic in all the specimens. They found ;
(a) for any given temperature the flake irons showed much greater growth than corresponding nodular irons;
(b) at the lower temperature an incubation period is required for initiating growth both in flake iron and in nodular iron;
(c) during the structural breakdown the rate of growth increases with in-creasing temperature;
(d) after completion of the structural breakdown flake iron continues to grow with time and temperature. The authors explain this phenomen by the assumption that continued growth is due to oxidation.
Ensor''^^' tested several low-alloy flake and nodular cast-irons in the 5 5 0 ^ - 700°C temperature range, and Richards'^' did the same thing in the ran-ge 600° — 920°C. It was found that at subcritical temperatures growth increases with time and increasing temperature. A Si content above 4.5% (in the presence of about 2% Cr and Al) reduces the growth. The growth of nodular iron is also reduced by the addition of 0.5 - 1.9% copper, but flake iron does not exhibit a si-milar property, in general, growth after a given heat treatment was found to be greater in flake iron than in nodular iron.
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Pugh tested the influence of carbide-stabilizing elements and a low-carbon equivalent, at 800^C and during repeated heating and cooling cycles. He maintains that growth is produced by internal oxidation and by the structural breakdown of the cast-iron, transforming pearlite and graphite. He concludes that the chemical composition of the iron is also an influential factor: as the silicon content increases (up to about 3.5%) the rate of growth also increases, due to the fact that silicon promotes graphitization. A silicon content above 4%, however, retards growth. Carbide stabilizing elements such as Cr, Mo and V, will produce a marked reduction in growth as well as a low carbon equivalent value by increasing carbide stability.
Grant"^'investigated the influence of V and Cr additions on pearlite stability. For this purpose he measured the changes in tensile strength, hard-ness and combined carbon, as functions of heat treatment at 600°C and 675°C. He established that increasing the quantity of V added and alloying by Cr, lowered the change rate of the above-mentioned properties.
In another investigation Grant'^' studied the growth characteristics of ingot mould iron. He suggests that in the vicinity of the critical temperature growth occurs thanks to the decomposition of carbides and to the oxidation of the metallic matrix. According to his calculations (based on the specific gravi-ties of ferrite, carbide, and pearlite. of between 7.5 and 7.9, and of graphite — 2.3) the precipitation of 1 per cent of graphite causes a volume increase of ap-proximately 2 per cent or a linear increase of 0.65 per cent, or 0.68 per cent ac-cording to some other authors'^^'.
To explain oxidation and cracking phenomena he cites several authors as having concluded that the main cause of growth is the oxidation of the silicon dissolved in the iron, with bulky silica thus being formed. This silica formation, together with the possible oxidation of iron and graphite, causes different rates of volume increases, depending on subsidiary conditions. The authors cited also showed that in gray iron having silicon contents varying from 1 to 6 per cent, the total growth during repeated heating to 900°C followed by cooling was approxi-mately proportional to the silicon content. This at least partially contradicts the observations of Pugh'^" and other investigator, who found a high silicon con-tent (4 — 6 per cent) retards growth (Silal type iron).
Grant established that during heating, especially cyclic heating in the
a t y transformation region, cracking takes place and that this also accelerates
oxidation. Grant cites suggestions to the effect that, upon heating, gas is liber-ated, probably at the boundaries between the graphite flakes and the iron, and the rapid expansion of that gas may give rise to an internal pressure sufficient to deform the metal permanently. It is conceivable that cracks or cavities, caused by the evolution of gas under pressure, can be found. It is very remarkable that in the cited literature the investigators obtained such widely divergent and even con-tradictory results from measuring growth in various atmospheres — comparisons being made with equal atmospheres — in approximately the same temperature range (repeated heating 850°— 900°C). For example, in some cases of vacuum treatment increased growth was found to take place, in others — just the opposite. Grant himself established in his own experiments involving repeated heating to 850°— 900°C (using material melted in vacuum and treated in air or vacuum), that growth can occur due to causes other than graphitization and oxidation. He reported that he did not reach a stage at which growth ceased; that the presence of gases dissolved in the iron or in the surrounding medium is unnecessary for promoting growth (as proved by melting the material with the highest growth and allowing it to solidify in vacuum). By previously annealing in air, and permitting oxides to form in the iron. Grant succeeded in inhibiting growth, but not indefi-nitely. Annealing in air at 500°C and 700°C resulted in some spheroidization and decomposition of the pearlite. The graphite flakes near the edges of the specimens
were surrounded by a dark layer of compact oxides and by an outer envelope, which was probably due to the oxidation of the silicon present in the ferrite to silica or silicates in a very fine form.
Growth (and scalling) of cast-iron at temperatures below and above the crit-ical temperature are surveyed by Hughes'^^'. He notes in his survey that early investigators thought gases in cast-iron to be the cause of growth but that this theory has received little or no support in recent years.
On the strength of investigations by others, it is shown that at subcntical temperatures, growth due to graphitization ceases after a certain time and a con-stant volume is maintained, but growth in air produces results invariably in ex-cess of the calculated growth. Referring to several authors, Hughes suggests that relatively small amounts of certain alloying elements are able to delay, or even to prevent altogether the graphitization of pearlite, i.e. they are also capa-ble of retarding or inhibiting growth. Another possibility is to raise the silicon content to such an extent that a ferritic as-cast structure is obtained. Yet an-other method is the administration of some prior heat treatment for achieving full graphitization.
It is well-known that white irons, which are virtually free of graphite, do not grow at subcritical temperatures, (since obviously no graphitization occurs).
Heselwood and Pickering''^' tested a series of ingot mould irons in low vacuum at several subcritical temperatures and with cyclic heating through the critical range. In vacuum the maximum growth was 0.5 per cent, obtained at 700'C during 5 hours, in air growth was slightly greater (about 0.6 per cent). For calculating the growth taking place during 5 hours at 700°C and other condi-tions established by them, these authors propose the following formula;
G = 0 . 3 2 - 0 . 2 1 Mn%+0.13 S i % - 0 . 7 3 P %
Correspondence between measured values and those calculated by means of the formula was only limited.
dissociated into a-iron and graphite, and that the carbon then migrates through the ferrite matrix, adding "secondary" layers to the already existing graphite flakes. This migration starts from the carbides nearest to the graphite flakes. As the graphitization progresses, so the thickness of ferrite across which the carbon atoms have to diffuse increases, the rates of graphitization and of growth accordingly decrease, and after a certain time, little or no further growth takes place. This model is true for subcritical temperatures. As a rule it i s suggested that the progress of graphitization will depend upon the mobility of the carbon atoms and on the solubility of carbon in ferrite, so that maximum growth will be obtained at the temperature at which the solubility of carbon in a-iron is highest.
Many investigators have stated that the most important factor in the growth mechanism is the decomposition of the cementite in the pearlitic phase. The dissolution rate of cementite (the graphitization rate) has been measured by sev-eral authors. They used the change of hardness (or other mechanical properties) in several types of cast-iron in conjunction with metallographic examinations'^^''^'^'
The relevant tests were carried out isothermally, at subcritical tempera-tures, or by cyclic heating. It is interesting to note that in some cases after a certain time, the change in the mechanical properties did not follow the general rule.
Nagaoka'^*' offers a very original and rather intriguing explanation for the cause of growth. According to his observation, an oxidizing atmosphere promotes
growth at austenitic temperatures, while at the critical temperature it has a re-tarding effect. He found that growth of cast-iron can take place either without graphitization of pearlite or without the oxidation of the matrix. He holds that while the cast-iron is being heated in the austenitic range, part of the graphite will dissolve in the austenite thanks to its increasing solubility, thereby produc-ing a "cast-off" void. Durproduc-ing coolproduc-ing, on the other hand, carbon solubility is reduced and the graphite precipitates, not necessarily into the "cast-off" voids but at least partly elsewhere, so that after the heating cycle some of these voids are left unfilled. This results in growth. Improving on this explanation the au-thor points to the behaviour of white cast-iron, where only cementite dissolves in the austenite, without any voids occuring; there is thus no growth, unless
graphitization has taken place.
He suggests a similar explanation for growth at lower temperatures. Ac-cording to his theory, the irreversible migration of graphite again results in the formation of cast-off voids (the author does not enlarge on this mechanism); but in contrast to the precipitation of temper carbon from austenite, the graphitiza-tion of the pearlite i s retarded, probably due to oxidagraphitiza-tion during repeated ing, so that the volume of undecomposed pearlite increases with repeated heat-ing. The growth rate thus gradually decreases.
Gilbert'^^' studied the growth properties of flake graphite cast-irons at 400°C and 500°C after several pre-heat treatments. He concluded that growth takes place most readily when oxidation of the matrix occurs, and that heating through the critical temperature range increases the susceptibility of the iron to oxidation. Still according to Gilbert, heating several times through the critical temperature range should increase oxidation and growth a these temperatures. In another work Gilbert'^*' observed the growth properties of several kinds of cast-irons during 6 years at 350°C and 400°C. There was practically no growth at 350°C. At 400°C less growth was found in nodular iron than in flake graphite iron of a similar matrix structure.
It will appear from the survey of contemporary literary sources that the rate of growth is possibly influenced, and perhaps even controlled, by several processes, viz.
(a) decomposition of carbides;
(b) oxidation by absorbed or penetrated gases; (c) cracking or bursting;
(d) diffusion of carbon towards graphite and of iron away from the surface of graphite.
It may be concluded that at this juncture there is no theory that accounts satisfactorily for all the phenomena associated with the growth of cast-iron''^'". especially at subcritical temperatures, and there are even conflicting opinions
as to the principal factor and whether all of the four phenomens enumerated ac-tually exist and, if they do, whether they play any significant part.
There seems to be doubt that if decomposition of carbides takes place in cast-iron, this must cause growth due to the formation of the more bulky graphite. The volume increase from this cause is about 2 per cent (which includes 0.68 per cent increase in length) for the formation of 1 per cent graphite.
The oxidation theory is less clear-cut. There are differences between the observations of several investigators with respect to the effect of oxidation, es-pecially at sub-critical temperatures. The influence of the absorbed or penetrat-ed gases is very uncertain. The experiments on the basis of which gases are claimed to be responsible for growth are not very convincing, as the observed growth might easily have been due to other causes, and there is in any case no information concerning the gases involved.
The cracking and bursting theory is very convenient for explaining growth
(30)
but has the disadvantage that the alleged cracking cannot be visually observed The carbon and iron diffusion theory as described in the works already mentioned appears to provide a reasonably good explanation, especially for the situation in the austenite range, but it is too incomplete to account for behaviour at subcritical temperatures, particularly in isothermal heating.
To sum up this discussion of previous works, it would appear that the var-ious hypotheses proposed succeed only partly in marshalling all the factors re-sponsible for growth at subcritical temperatures.
It should be noted in this connection that a comparison of the results ob-tained by the various authors is very difficult, for the following reasons;
1. As a rule, details are given of the composition of the cast-iron investigated as well as of the conditions in which it was pro-duced, but for gaining a fully informative picture a special re-search would be required to compare the separate and the com-bined influences of the composition, the melting and pouring temperatures, the atmosphere, the mould materials, the cooling
rate. etc.
2. In most cases the as-cast structure is given, and this alone already points to differences between the various experimental materials. Moreover, many authors administer some advance heat treatment to their specimens in order to obtain uniformity of structure before embarking on the actual growth experiments. Needless to say, such treatment is apt to introduce differences that later affect growth behaviour.
3. Yet other differences exist among the various methods and cir-cumstances for measuring growth, such as specimen size, tem-perature. heat control, and ambient atmosphere. When this latter
item is named, it is often hard to establish whether the gas used was of technical or laboratory purity, or in the case of vacuum
measurements, what degree of vacuum was maintained.
4. Several authors have measured growth and graphitization (phase change) during successive interruptions of the heat treatment given. The data measured were changes of dimensions, hardness or other mechanical properties, and changes in the micro-structure"^"'^-^^'. Others - and the present author amongst them — measured growth dilatometrically, i.e. without interrupt-ing the heat treatment. Obviously, this precludes a truly com-parative examination of the microstructures.
It is accordingly understood that experiments made under conditions that are only ostensibly similar produce widely dissimilar results.
2.4 P R E V I O U S O W N W O R K O N T H E S U B J E C T
Prior to the present work an investigation was carried out into the factors affecting the service life of small gray cast-iron ingot moulds as used in a cer-tain steel mill'^'. In these moulds, which weighed approx. 800 kg and had a wall-thickness of about 60 ram. carbon steel ingots of 450 kg were cast at a
tempera-tare of 1600°- 1620°C.
In the course of that investigation the thermal gradients caused in the mould during heating up and cooling down were determined. Another object was to study the changes in the microstructure of the cast-iron during the service life of these moulds.
Due to the practice of bottom pouring the maximum temperature in the mould wall was reached at the bottom, about 20 — 25 minutes after the beginning of pour-ing. That temperature never exceeds 700°C at a distance of 5 mm from the work-ing face. All other temperatures varied, reachwork-ing up to 500°C dependwork-ing on the location of the measurement. On the working face of the ingot moulds cracks ap-peared after a number of pouring cycles, caused by temperatures higher than the critical temperature as well as by oxidation. These cracks propagated during service-life of the mould and adversely affected the'surface quality of the ingots; but only very rarely at any rate, in the course of the investigation, could fractur-ing of the moulds be attributed to them. Most of the fractures examined had their origin in the external face of the moulds. The conclusion was accordingly reach-ed that the decisive factor determining the service-life of the moulds were the in-ternal stresses caused by the variations in the phase change rates within the
mould walls, such variations being, of course, due to the thermal gradient. Another work which preceded the present research dealt with the method to be used in the investigation. At first it was suggested to study the behaviour of cast-iron at subcritical temperatures in conditions of cyclic heating. A device was constructed which introduced samples into a 700°C salt bath for 30 minutes and subsequently withdrew them for air-cooling during another 30 minutes. This was intended to imitate the conditions of cyclic heating prevailing inside an in-got mould. The changes in the cast-iron were studied by means of micro-struc-tural changes and the variations in hardness.
An analysis of the experimental results led to the conclusion that the method described is suitable only for qualitative comparisons among several kinds of cast-iron, and that it is unable to provide a basis for the quantitative assessment of the kinetics of the different changes. Another conclusion worth
to mention is that it appears to be virtually impossible to arrive at an objective quantitative estimate of the amounts of the different phases present in the micro-structure, especially in the case of flake graphite. Hundreds of volume measure-ments, carried out in this preliminary work with several methods (linear and point counting, "Quantimet"). proved that the standard deviation is so large that the mean value obtainable is useless for the assessment of the composition of the structure.
This conclusion appeared to be inevitable, although it contradicted those reached by other investigators'^"'. The explanation is to be sought in the heter-ogeneity of the cast-iron. It will be readily appreciated that the cementite phase present in the pearlitic matrix can be measured only through very high magnifica-tions at which, however, the amounts of cementite may vary between zero and the maximum from the one site to the next. Moreover, due to the random distribu-tion of the graphite the volume of the latter may also vary in a like manner. At low magnification, on the other hand, the graphite boundaries are not easily dis-cernible and thereby, too, the measurement becomes subjective and therefore in-accurate. Yet another reason for the difficulties attending the determination of the graphite volume will become clear in the course of the present research.
In the light of the foregoing it became obvious that other methods had to be devised for investigating the behaviour of cast-iron at subcritical tempera-tures. It was found that the behaviour of cast-iron during isothermal heating was the same as during cyclic heating to a peak temperature equal to that of the iso-thermal heating. Isoiso-thermal heating at a high subcritical temperature was there-fore decided upon for the present study.
For observing the most fundamental changes, dilatometric measrurements were chosen, since they appeared best to guarantee an objective assessment. Subsidiary experiments served to explain the basic phenomena indicated by dila-tometry.
3. E X P E R I M E N T A L W O R K A N D M E A S U R E M E N T S
3 . 1 K I N D S O F C A S T . I R O N U S E D IN T H E E X P E R I M E N T S
Four series of flake-graphite cast-irons, each of different compositions. were prepared for the experiments, and parallel series of nodular cast-irons hav-ing roughly the same compositions were also cast with a view to enablhav-ing com-parisons to be made, if such should be found to be necessary.
All these irons were melted down in an induction furnace in batches of 80 kgs. Synthetic pig iron and low-carbon steel were used as raw materials, the
re-quisite alloying elements and inoculants being added wherever necessary. The eight series were melted in four batches.
The necessary silicon content was obtained by the addition of ferro-silicon to the melt while in the furnace and by calcium silicon ladle additions. All other elements were given as ladle additions. Inoculations for obtaining nodular types of cast-iron consisted in addition of FeMgSi and calcium silicon. All the melts were superheated to between 1500° - 1550°C and poured into green-sand moulds at 1350°- 1360°C, a keel-block shape as shown in fig. 3.1 being chosen. :a)eci-mens for the experiments were taken from the centre of the keel. For the composi-tions of the eight cast-irons see table 3.1. The microstructtu-al examination show-ed all the series of the testshow-ed flake-graphite iron to be of as-cast pearlitic struc-ture with flake graphite Type A, Size 4, according to the ASTM A-247 classifica-tion. All tests were made with as-cast material, no pretreatments whatsoever being given. 1 60
V
- - 11,
r 160 155 F I G . 3. 1 K E E L S L O C KT A B L E 3.1
COMPOSITION OF T H E CAST IRON USED.
C o Si % Mn % Ct% C u % P% S% H—2 Flake graphite cast
H - 3 F l a k e graphite cast H - 5 Flake graphite cast H - 6 F l a k e grsiphite c a s t H-8 Nodular c a s t iron H-9 Nodular c a s t iron H-11 Nodular c a s t iron H-12 Nodular c a s t iron iron iron iron iron 2.96 2.90 3.00 2.94 3.00 2.95 3.04 3.00 2.07 1.97 1.82 1.94 1.81 1.91 1.77 1.72 0.52 1.15 0.52 0.51 0.47 1.15 0.44 0.44 0.021 0.017 0.017 0.020 0.020 0.020 0.015 0.018 0.012 0.012 0.012 0.012 0.014 0.014 0.013 0.010 3.2 D I L A T O M E T R I C M E A S U R E M E N T S
In order to establish the growth kinetics, the cast irons were subjected to dilatometric measurements. The specimens for the purpose were 50 mm long and and of 4 mm dia. The measurements were effected by means of a transducer ac-tuated by a silica rod connected to the free end of the specimen and thus trans-mitting the dilatation of the specimen.
The specimens, enclosed in a silica tube, were heated in a high-purity ar-gon atmosphere.
The tests were performed for 165 hours at 700°C. This relatively high subcritical temperature was chosen for the isothermal heating because prelimina-ry experiments, made with all the kinds of cast iron used, had shown that the A, critical temperature was situated within the range of 730° — 760°C.
The thermal treatment was continued until the dilatation curve reached a constant level, signifying that no appreciable further growth was to be expected.
curves, and hence the curve for iron H—6 is not shown in figs. 3 2 and 3.3. which represent the incremental dilatation AL as a function of time. The in-cremental dilatation is in all cases proportional to the transformed fraction. Fig. 3.2 shows the dilatation curves as obtained originally, while fig. 3.3 shows di-latation as a function of the logarithm of time.
500 400 300 200 100 50 0 ^\(n) 1.000 ^ ^
/f\^
/ 5.000 _ _ _ - H_2 H - 3 ^ _ _ _ _ ^ _ _ — - H_5 10.000 . mm F I G 3.2 OIL A T A T I O N C U R V E S O F F L A K E G R A P H I T E C A S T I RONS F I G . 3 3 D I L A T A T I O N C U R V E S O F C A S T - I R O N S H - 2 H - 3 A N D H - 5 l l m . n , AS F U N C T I O N S O F T H E L O G A R I T H M O F T I M EThe analysis of these sigmoidal curves shows that they follow the equation (3.1) y = 1 _ e ^^
In this equation y — tne fraction of growth in time t.
n — a constant, referred to as time exponent index, K — temperature-dependent rate constant.
This being so, the graph of loglog vs. log t must be a straight line, 1 - y
from the slope of which the value of n is obtained. The graph mentioned is shown in fig. 3.4 . The value of K is deduced either from the intercept or by the solution of eq. (3.1). K may also be obtained from the data of the transformation curves; rearrangement of eq. (3.1) shows that — equals t", the time at which
K y - S-=^ = 0.6321
A computer program was drawn up for checking the graph and for exactly -^ determining the values of n and K. For that program, eq. (3.1) took the form
(3.2) AL 1 - e - K t '
where AL — linear expansion at 700°C, measured dilatometrical-ly as a function of time, and
Y„ — Final expansion as t -. 0 , i.e. the transformed
r
fi-action equals 100% (or AL = 1)
The curves obtained by means of the computer fit those obtained experi-mentally with only minor deviations, the largest of which occiu's during the first 10 hours. It is of interest to note that Y„ differs only very slightly from the
r
expansion taking place after 165 hours (see Figs. 3.2, 3.3). The data obtained from the computer can be summarized as follows;
n K Y p W H - 2 ( H - 6 ) 1.16 1.3 X 1 0 - " 476 H - 3 1.74 9.81 X 1 0 - ^ 457 H - 5 1.29 2.64 X 1 0 -456
In order to enable comparisons to be made, the dilatation experiments were carried out with the nodular ^ e c i m e n s . The results of these experiments are shown in fig. 3.5. It will be seen from them that the maximum dilatation of H—8 nodular cast iron is about half that of the H—2 flake graphite iron which is of a similar chemical composition, while the time required for reaching it is about a third. H—9, H—11 and H—12 nodular cast irons had free cementite in the as-cast structure, and this proved impossible to decompose at 700°C, as will be discussed in section 8.1.1.
Mill) 300 250 200 150 100 50 0 / / / / ^ = ^ 500 ' / / / / ^ 1.000 - ^
X
^ 1.500 ^ ^ ^ ^ 2.000 — . H _ n __-^ ' H - l l / ^ : ^ 2.500 3.000 -^ H-12 , . - - ' ' ' ' ' " t i m e - m i n . 3.500 4.000 4.500 5.000 F I G . 3.5 D I L A T A T I O N C U R V E S O F N O D U L A R C A S T - I R O N S 3.3 D I F F U S I O N C O U P L E SThe pearlite decomposition being described in this work in terms of carbon migration through ferrite as a medium, the influence of the elements dissolved in the latter was investigated. For this purpose, diffusion couples were prepared'^^', one side of which was constituted in turn by each of the four types of flake graph-ite cast iron (Table 3.1), while the other side was chosen as follows;
Series a — pure alpha-iron,
Series b — Alpha-iron containing Si, Series c — Alpha-iron containing Cr,
Series d — Alpha-iron containing Cr and Si. For the exact compositions .if the alpha-irons see Table 3.2.
T A B L E 3 . 2
C O M P O S I T I O N S O F T H E A L P H A - I R O N S U S E D
C % Si % Mn % Cr %
1. Swedish Pure Iron 0.02 0.005 " 0.005 0-005 max.
2. Swedish Iron -I- Cr Alloy 0-1 3. Swedish Iron -I- si Alloy 1,42
4. Swedish iron+Si+Cr Alloy 1.78 0.18 A typical diffusion couple used in this study is shown in Fig. 3.6. Before
the joint, the interfaces were thoroughly cleaned and polished. During the heat treatment the parts of the couple were tightly locked together by means of a cast-iron nut. Since at high temperatures the expansion of the a-phase is greater than that of the cast iron, the tightening increases with temperature. The diffu-sion couples were enclosed in a silica tube, which was evacuated sealed her-metically, and placed in a temperature-controlled furnace in order to keep the specimens at a constant temperature viz. 700''C for 72 hours. After this heat treatment the specimens were left in the furnace to cool.
F I G . 3.6 T Y P I C A L L O C K I N G O F D I F F U S I O N C O U P L E ; G R E Y CAST I R O N WITH A L P H A I R O N . M A G N . 4 X
Since the experiments showed that results from H—2 and H—6 cast-irons were practically identical, as were those from H-3 and H - 5 . data are given only for the H-2 and H—5 types respectively although, of course, actual experiments were made with all four types throughout.
In the series of couples joining unalloyed cast iron (No. H—2 in Table 3.1) with unalloyed alpha-iron a large amount of cementite was found to have formed on the alpha-iron side of the interface (fig. 3.7). As can be seen from the figure, the carbides constitute very regular globules along the interface. In the couple chromium-alloyed cast-iron with pure alpha-iron (No. H—5 in Table 3.1), the amount of cementite observed appeared to be smaller (fig. 3.8).
In the second series of diffusion couples the cast-irons of the first series were joined to alpha-iron containing Si (No. 3 in Table 3.2). In none of these couples, regardless of the alloying elements, could any traces of cementite be detected at the interfaces (figs. 3.9, 3.10).
The third series of couples were made with samples of the cast-irons list-ed in Table 3.1 joinlist-ed to chromium-alloylist-ed alpha-iron (No. 2 in Table 3.2). In the couples with unalloyed cast-iron, a much larger amount of cementite (compar-ed with the first series) was found on the alpha-iron side of the interface, but none in the couples with chromium-alloyed cast-iron (figs. 3.11, 3.12.).
In series " d " — the different cast-irons joined to alpha-irons containing both Si and Cr (No. 4 in Table 3. 2) — no cementite whatever formed at any of the interfaces except in the couple with unalloyed cast-iron, and that in very small quantities (fig. 3.13).
^^ V
>N-V-:^/:
- - v ^ l
J •w "^^s .^V/ • •FIG. 3.7 INTERFACE OF A DIFFUSION COUPLE OF H-2 CAST IRON WITH ALPHA IRON. UNETCHED, MAGN. 150 X
^ V
• .— ^ /
FIG. 3.8 I N T E R F A C E OF A DIFFUSION C O U P L E O F H-5 CAST IRON WITH ALPHA IRON. UNETCHED, MAGN, ISO X
F I G . 3.9 I N T E R F A C E O F A D I F F U S I O N C O U P L E O F H - 2 C A S T I R O N W I T H SI C O N T A I N I N G A L P H A I R O N . UN E T C H E D , M A G N . ISO X /
' ^ ^l ^N4-'^1
-r^y-- /V ^
( 'L
. i/
F I G . 3.10 I N T E R F A C E O F A D I F F U S I O N C O U P L E O F H - B C A S T I R O N WITH Si C O N T A I N I N G A L P H A I R O N . U N E T C H E D , M A G N . ISO X^ ' ^ J a ^
. V |
FIG. 3.11 I N T E R F A C E O F A DIFFUSION COUPLE OF H-2 CAST IRON WITH CHROMIUM CONTAINING ALPHA IRON. UNETCHED, MAGN. 150X
FIG 3.12 I N T E R F A C E OF A DI FFUSION COUPLE OF h - 5 CAST IRON WITH CHROMIUM CONTAINING ALPHA IRON. UNETCHED. MAGN. 150 X
F I G . 3 . 1 3 I N T E R F A C E O F A D I F F U S I O N C O U P L E O F H - 2 C A S T I R O N W I T H SI A N D C r C O N T A I N I N G A L P H A I R O N . S P E C - E T C H I N G , M A G N . 300 X
alpha-iron side of the interface, starting from the decomposing pearlitic matrix of the cast-iron and from the graphite lamellae.
Microprobe analyses were made in order to check whether, in addition to C, any of the original cast-iron elements had diffused into the alpha-iron. In all the couples studied here the results were negative.
Some typical results of the electron microprobe analyses are shown in figs. 3.14, 3.15, of which figs. 3.14 (a) (b) (c) (d) (e) refer specifically to the carbon dispersion and figs. 3.15 (a) (b) (c) (d) to the dispersion of other elements.
3.4 D E N S I T Y M E A S U R E M E N T S
The densities of all the kinds of cast-iron used in this investigation were measured in the as-cast state and after the completion of the heat treatments.
FIG. 3.14
H-2 CAST IRON + FERRITE C SCANNING LINE
(a) (b)
H-6 CAST IRON + FERRITE C SCANNING LINE
(c) (d)
> •
\ - ^
H-2 CAST IRON -I-FERRITE WITH Si C SCANNING LINE
(e)
H-2 CAST IRON + FERRITE F e SCANNING L I N E ( a )
H-3 CAST IRON + FERRITE Mn SCANNING L I N E ( b )
^ /
-.i-"'
/ H-5 CAST IRON + FERRITE
Cr SCANNING L I N E ( c )
H-6 CAST IRON + FERRITE C u SCANNING L I N E ( d )
i.e, when all the carbon of the cementite had been transformed to graphite. The following measurements were made:
(a) Weight of the specimen in air;
(b) Weight of the specimen in distiled water;
(c) Weight of the wire holding the specimen during weighings (a) and (b). These were used in the following calculations:
(d) Weight of the specimens minus the wire from which it was suspended; (e) Volume of the specimen as the difference between its weights in air
and in distiled water, both with the weight of the suspension wire taken into account;
(f) Density of the specimen, being the quotient of (d) and (e). The results were as follows :
In the as-cast state the density of the flake-graphite iron was 7.40 g/cm^ — 7.48 g/cm^. The density of the same type of iron after heat treatment ranged from 7.16 g/cm' to 7.25 g/cm^.
The density of the nodular irons in the as-cast state was 7.38 g/cm^ to 7.45 g/cm^, and after the heat treatments 7.25 g/cm^ - 7.30 g/cm'. .
3,5 HARDNESS CHANGES
Although it appeared to be evident, both from the analysis of the results obtained by several authors and from those produced in the present author's own previous studies, that the most suitable method for tracking phase changes during isothermal heating at subcritical temperatures is that employing and analyzing dilatometric measurements, it was felt advisable to represent these phase changes by their effect on the hardness, too.
For the requisite hardness measurements specimens were prepared from the different cast-irons listed in Table 3.1, placed in a protective atmosphere, and
heat-treated at 700°C for those periods which were indicated by the dilatometric measurements as marking the ends of the various phase changes, i.e. when dila-tation temporarily ceased. After each such period certain specimens were withdrawn, air-cooled, and subjected to hardness tests. The different hardnesses measur-ed were plottmeasur-ed as functions of the heat-treatment periods. Fig. 3.16 shows the results. .- N r4 n n « a a — — 1 1 1 1 1 1 1 1 I I I I I I I I • • « a « o « i B +
niMi r^:^
2 A 1 1 I 1 I1
; 1 < * 4 t 1 ( 1 1 1 1 1' V
/ 0 ' 7 ? / •/ ,•' < / • , ' • / 4 / X / / y^ • , ' ' .•'* •* ^ ^ ° "E cb " o o "" ^ o t o . ^ o ^(N IN (N cc u 3 m ^o CN o «^ » 1 t . ^ 1 I I I ^ . . « CD ••-0 1 0 . ^ * • • . . \ * * : : i tj
J •s 1 : t 0 < • : •; < .' /• < ° / > < // n
' .' ••• / * f - i / ' • ' / // /, yy'y* •
""^ o O 1—1 *.* o • » ""s
• " -O t— o c o 0 0 o r-^ S o o IT) o >»s
o <N 5 o \ri o • — • "3.6 M I C R O S T R U C T U R A L CHANGES
The specimens prepared for the hardness tests as described in Section 3.5 above were also used for microscopic examinations. Each specimen wsa photo-graphed in the as-cast state and again after each heat-treatment period.
At a certain stage of the experiments it was proposed to utilize the same specimens for the quantitative determination of the progress of the volume frac-tions of those phases that are subject to changes, viz. cementite and graphite. However, after hundreds of measurements were taken on several specimens, using various methods ("Quantimet", point counting, and linear counting), it became ap-parent that due to the great heterogeneity of cast-iron no reliable quantitative
as-sessment of the different phases is possible. More specifically, measurement of the eutectoid cementite particles is possible only at very high magnifications, and in these conditions the cementite quantities determined may vary between zero and quite unrealistically high values. This is particularly true when, as a result of proponged heat treatments, some of the cementite has already dissolved. The mean value calculated is extremely doubtful, a fact that is borne out by the stand-ard deviation value even after over one hundred measurements were taken from one and the same specimen. As to the graphite measurements, similar considera-tions apply, although here lower magnificaconsidera-tions seem applicable. In that case, however, the determination of the boundaries, particularly of flake graphite, is subjective to such an extent that the inaccuracy of the measurement exceeds the expected volume change from one specimen to the other. In order to give an idea of the morphology of the phases dealt with, figs. 3.17 (a) (b) (c) (d), 3.18 (a) (b) (c) (d) are presented.
H - 5 M A G N 1 0 0 0 : ( c ) AS CAST S T R U C T U R E S OF T H E F L A K E FIG. 3.17 H - 6 M A G N lOOOx (d)
t
G R A P H I T E CAST IRONSH - 5 32 H O U R S . M A G N . 1 0 0 0 x
(c)
MICRO-STRUCTURES AFTER SEVERAL H£
FIG. 3.18 H - 6 44 H O U R S . M A G N . lOOOx (d) * > *. T TREATMENT TIME AT 700° C
passing from the source to the target. A thin Na(Tl) crystal (0.1 mm) was mounted on a low-noise-level photo-multiplier RCA 8575. The pulses obtained from the scintillation counter were amplified and passed through a single-channel analyzer to the Elbit-100 computer working as a multiscaler analyzer.
3.7.2 P R E P A R A T I O N OF SPECIMENS FOR MOSSBAUER MEASUREMENTS
The specimens were prepared from H—2 cast iron (see Table 3.1), in the as-cast state and after different periods of heat treatment at 700°C.
For purposes of comparison a special specimen was cast in vacuum from high-purity Fe and high-purity Si, the silicon content in the final alloy being
1.68%. Carbon content was 0.005%.
The results are shown in figs. 3.19; 3.20; 3.21; 3.22; 3.23; 3.34; 3.25.
3.7.3 I N T E R P R E T A T I O N OF THE MOSSBAUER SPECTRA
Fig. 3.19 shows a typical Mcissbauer spectrum of the H-2 cast iron, while fig. 3.20 pertains to F e ^ c . Satellites are seen to be present in addition to the alpha-iron peaks in the spectrum of the cast irons.
Only the first peaks (the lowest velocity) are shown and will be the only ones to be considered. Figs. 3.19 and 3.21 show that the first satellite may be attributed to a Fe^^ atom adjacent to a Si atom, while the second satellite — as is demonstrated in Fig. 3.19 derives from a Fe^" atom adjacent to an inter-stitial carbon atom. These atom sites are in an octahedral interinter-stitial position similar to the situation existing in Fe-C martensite, in both which cases distor-tion dipoles result. This will be discussed in greater detail in Secdistor-tion 7.3.1. The point of interest to note in this connection is that approximately the same satellites, namely the first and the second appear in a quenched silicon steel
98 1 ° 97 W) H ^ 9 6 o u 95 94 93 92 91 90 r '""^ ^ \ / ^ / /SECOND / SATIELLITE
1
I 1J
1 FIRST /SATELLITE1/
_0._S? - 0 - 4 - 0 . 3 - 0 . 2 r 1 1 300 400 . . ' fc*. * - 0 . 1 0 1 500 . - • ^ ,-0.1 1 " ^- .-' 0.2 0.3 1 600 ", -0.4 • 0.52 , ^ „ ^ , , ^ , 1 1 700 800 A D D R E S S CO F I G . 3 . 1 9 M O S S B A U E R S P E C T R U M O F H - 2 IN A S - C A S T S T A T E49 I o 69 68 67 66 -65 H 64 M 6 3 i 62 u 61
I ' l l
y
/i!
S
I
i
• • ^ ' ' ' ' ' • • r . . • • 11 • • • I f . . . . . 1 1 r> inn 1 nn inn .,.«>. _ -200 300 4 0 0 500 6 0 0 700 ADDRESSFIG. 3.20 MOSSBAUER'S SPECTRUM OF FE^C
(0 38% C, 1.2% Si) - Fig. 3.26*^^' - and in cast-iron in the as-cast state — Fig. 3.19 . It may therefore be said that thanks to the silicon dissolved in ferrite the amount of dissolved carbon increases and that more or less the same effect is obtained in cast iron without quenching, as in quenched steel.
In Figs. 3.22; 3.23: 3.24; 3.25, a comparison is made between specimens made of H—2 cast iron slowly cooled after different heat treatment periods at 700-C. It will be seen that the second satellites decrease with heat treatment time, as do the first satellites. The decrease in the first satellites is due to the fact that as a result of cementite decomposition greater volumes of ferrite are generated, and the average Si content in the ferrite decreases in consequence. At the same time the decrease in the second satellites proves the author's point about the positive interaction between Si and C atoms. Therefore, the lower the amount of silicon dissolved in the ferrite. the lower also the amount of the
m ' o
1 **
1 3 o u 1000 j 985 975 950 925 . 900 - ^ • [r>5v
U
/J
W[
- 0 . 6V
- 0 . 5 300 / ^ < - 0 . 4 - 0 . 3 . , VELOCITY em/sac ' ADDRESS 40098. 97-96 -• 95 94 9 3 -SECOND SATELLITE FIRST SATELLITE 92 •• 91 •• VELOCITY cm/sec. -0.52 — I -0.4 -0.3 90 . . ADDRESS I 300 I • 400
98 •• - . - 97 • • 96 •• 95 •• 9 4 - • 93 •• FIRST SATELLITE 92 9 I -V E L O C I T Y c m / s e c . -0.52 -0.4 — I I — -0.3 — » - 90-f-A D D R E S S 1 300 I 400 F I G . 3.25 H - 2 C A S T I R O N A F T E R 10 H O U R S H E A T T R E A T M E N T
carbon dissolved in it. The slow cooling of the specimens in connection with the subcritical temperature was found to be advisable as a reliable means to en-sure high solubility of carbon in ferrite not as a result of quenching but due to the presence of Si.
c
c
c
; UJ -H . 1 U K ^ \tntn o 00 'i FIRS T If SATELLIT E 00 •o juno;;) -— ^ -* r^ M 111 K Q Q <-1
-§ o o o4. D I S C U S S I O N O F T H E G R O W T H K I N E T I C S
The dilatometric measurements confirm that, from a kinetic point of view, the growth of cast-iron at subcritical temperatures resembles that at temperatures above the critical point, it also bears a resemblance to the growth of pearlitic steel at subcritical temperatures. In both the latter cases investigators are unanimous that the equation:
y = 1 - e ~ ^ '
is suitable for representing graphitization kinetics, and the experiments carried out in the present work confirm this also for the growth of cast-iron at subcritic-al temperatures. This means that the principsubcritic-al factor governing growth kinetics is the normal heterogeneous process occuring by virtue of the growth of one or more phases at the expense of others, in other words, because:
a) the absolute value of the growth as a function of time never decreases (there is no contraction at any time),
b) growth does not continue indefinitely, the growth function approaching y = 1 asymptotically.
c) the growth rate is not constant.
In the case of growth at subcritical temperatures, it may therefore be stat-ed that the one process which can be^considerstat-ed as invariably governing growth is carbide decomposition attended by graphitization. None of the other proces-ses suggested appears to fulfil all the above conditions.
With the data presented in Section 3.2 to hand the growth process can now be analyzed mathematically. The derivative of eq. (3.1) with respect to time i s ; (4.1) ^ = n K t " - l e-Kt"
The maximum of the derivative curve indicates the point of inflexion of the sig-moidal curve. The derivative of eq. (4.1), equated to zero, yields —
(4..2) ^ = n K e - K t " [ _ n K t 2 ( " - l ) + ( n - l ) t " - 2 ] = 0 dt2
whence
(4.3) nKt" = n - 1
while the time at which the inflexion occurs,
(4.4) t, - ° • ' " ^ ^
\J-.
K
Eq. (4.4) gives the following values of t^ for the flake graphite irons used in this work (see Section 3.2):
''i(H-2) '^ '*^^ mins.; H(H-3) ~ ^^'^^ mins.; *i(H-5) " ^^^^ mins.
Substituting eq. (4.3) into eq. (3.1) gives the dilatation at the point of inflexion: 1-n
(4.5) yj = l - e "
yielding the following values of yj for the same types of cast-iron: y.(jj_2) = 0 . 1 3 ;
yi(H-3) = 0-35; ^KH-S) = 0-20 .
The points t. and y. may be considered as characteristics of the rate at which the dimensional changes vary. The process of dilatation described by eq. (3.1) may accordingly be represented as taking place in two stages, which are se-parated by the point of inflexion.
Since the growth was made to occur in an oxygen-free atmosphere, the dimensional growth must, in principle, be due to the carbon's leaving the cement-ite and forming graphcement-ite. In other words, dimensional growth is caused by the differences in density and volume between the original and the new phases. In practice, however, actual growth in the case of flake graphite exceeds that to be expected from the volume and density changes alone. The reasons for this ex-cess growth will be discussed later, in chapter 6.
Growth is the composite result of three subsidiary processes a) Cementite dissolution in ferrite ;
b) Carbon migration through the ferrite ;
c) Graphitization in the course of which carbon atoms pass from the ferrite crystals to the graphite.
The rates of these three processes differ from one another, and in certain condi-tions each may reach a maximum. The growth rate will obviously be determined by the subsidiary process having the lowest rate. Looking at the three proces-ses from a metallurgical point of view, the following statements can be made; 1) The rate of cementite dissolution is constant for unit time and unit
area with respect to a certain type of cast-iron but may be affected by alloying elements, as will be discussed later (Chapter 8).
