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The bainitic mechanism of austenite formation during rapid heating
PII:
Acta mater. Vol. 46, No. 16, pp. 5917±5927, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00210-9 1359-6454/98 $19.00 + 0.00
THE BAINITIC MECHANISM OF AUSTENITE FORMATION DURING RAPID HEATING W. J. KALUBA1{, R. TAILLARD2 and J. FOCT2 Laboratoire d'Applications en GeÂnie Industriel-Energie-Environnement, Universite du Littoral CoÃte d'Opal, 62968 Longuenesse, France and 2Laboratoire de MeÂtallurgie Physique, URA CNRS n8234, Universite de Lille 1, 59655 Villeneuve d'Ascq, France
1
(Received 19 March 1998; accepted 7 May 1998) AbstractÐStrips of a spheroidized 0.68% C, 0.67% Mn and 0.24% Si pearlitic steel were submitted to rapid heat cycles. The samples were heated by direct electrical conduction up to a peak temperature and immediately water quenched. The morphological features of the a 4 g transformation were studied by light and electron microscopy. It was found that austenite at an early stage of the transformation has a speci®c lath-like morphology and that its growth always starts directly from a grain boundary of ferrite. Moreover, the presence of retained austenite allows the prediction of relatively high carbon content in the transformation product. The particular interaction of austenite laths with carbide particles gives an indication that the growth mechanism can be displacive at its early stage. Comparing the morphological characteristics of austenite formation observed in this study with those reported for austenite decomposition, a bainitic model of the early stage of a 4 g transformation is proposed. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. ReÂsumeÂÐLes premiers stades d'austeÂnitisation d'un acier perlitique globularise aÁ 0,68% C, 0,67% Mn et 0,24% Si sont eÂtudieÂs par chauage rapide. De ®nes bandes de meÂtal sont chaueÂes par eet Joule jusqu'aÁ une tempeÂrature maximale avant d'eÃtre immeÂdiatement trempeÂes aÁ l'eau. La morphologie du produit de la transformation a 4 g est eÂtudieÂe par microscopie optique et par microscopie eÂlectronique en transmission. En deÂbut de reÂaction, l'austeÂnite cro|Ãt directement aÁ partir d'un joint de grain ferritique et adopte une morphologie en lattes. La conservation de la nature austeÂnitique de ces lattes aÁ tempeÂrature ambiante suggeÁre son fort enrichissement en carbone. Les interactions observeÂes entre des lattes d'austeÂnite et des particules de carbure militent par ailleurs en faveur d'un meÂcanisme displacif en tout deÂbut de la transformation. Les eÂvolutions morphologiques observeÂes en cours de chauage dans cette eÂtude conduisent aÁ proposer un meÂcanisme de type bainitique pour la transformation a 4 g. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
In solids, the behaviour of the interface between parent and product phases is of chief importance in order to characterize any phase transformation. This behaviour is generally determined by the value of the driving force induced by thermal treatment, and also by the structure of the interface. For instance, an increasing driving force can cause a substantial change of transformation mechanism (from reconstructive to displacive), which usually goes with an evolution of the morphology of the product phase (from equiaxed to lath-like/plate-like in this case). One of the best illustrations of such a tendency is aorded by the widely studied iron±carbon system. The numerous papers on austenite decomposition discuss the relationships between the morphology of ferrite, the mechanism of the g 4 a transformation and the supercooling. The rich variety of resulting morphologies are classi®ed in the following well-known sequence: ferrite allotrio{To whom all correspondence should be addressed.
morphs±WidmanstaÈtten ferrite±bainitic ferrite and martensite at an increasing degree of supercooling [1±8]. Two major transformation mechanisms are reported to take part in the formation of one or other morphological state: a reconstructive one involving a time dependent, diffusional growth and a displacive one based on the movement of coordinated atoms occurring by shear. The reconstructive mechanism is generally accepted to control the growth of ferrite allotriomorphs at low supercooling, whereas the formation of martensite is proved to occur under high supercooling by the displacive manner. However, in spite of the number of works in this area, there are still controversial opinions on the growth mechanism of transformations that takes place at intermediate supercooling, especially concerning the formation of bainite [9±13]. The morphology and mechanism of reverse transformation from ferrite to austenite have been less widely studied, partly because of the number of experimental diculties encountered. The high temperature product phase is dicult to observe
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KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
directly without the use of sophisticated experimental techniques. Moreover, the morphological features of the transformation are usually modi®ed by cooling to room temperature. Consequently, the problems of austenite formation in the iron±carbon system are often discussed in terms of a well-known isothermal diusion-controlled growth, based on the local equilibrium concept and taking account of the dissolution of cementite in the matrix [14]. The kinetics of austenite formation during intercritical annealing was presented in several works [15, 16]. It was shown that the experimentally measured growth kinetics diers signi®cantly from the theoretical predictions for the temperature range of 740± 7808C and especially for short times of austenitization. It should be noted that the heating method used in these studies (in lead pots or salt bath) as well as the relatively small dimensions of the samples, enables us to predict a relatively high heating rate of the material prior to austenitization. Yet, the in¯uence of the heating rate on the a 4 g transformation at its early stage was not discussed. Moreover, the short times of austenitization (less than 5 s) were not taken into consideration. However, under conditions of continuous heating, especially at high rates, there must be some deviation from the equilibrium model. In close parallel with the well-known eects of cooling, the corresponding increase of the driving force should give rise to a change of the transformation mechanism and in consequence to a morphological evolution of austenite with heating parameters (peak temperature, heating rate . . . ). The problem was the subject of several studies [17±21], including the role of heating rate and the initial microstructure of Fe± C alloys. It was shown that an increasing heating rate brings about a continuous rise of Ac1 and Ac3 temperatures. Moreover, a distinct temperature plateau was observed at a heating rate exceeding some critical value. The occurrence of this plateau was sometimes considered as an indication of a change of the mechanism of transformation from reconstructive to displacive. Unfortunately, the amount of morphological evidence to support such a statement is rather limited, especially for steels with high carbon content. Zerwekh and Wayman [22], using pure iron whiskers, were the ®rst to prove that the a 4 g transformation can occur by a displacive mode even at low overheating. They observed a well-marked shape change during the transformation and de®ned the habit plane. They were not able, however, to observe the same feature on the polycrystalline material. With iron±carbon alloys dierent austenitization behaviours were noticed depending on their initial microstructure. It was suggested that the transformation starts at cementite±ferrite interfaces in a reconstructive way, whereas massive or displacive modes are able to operate at ferrite grain
boundaries [17, 23]. Displacive formation of austenite was later con®rmed in laser heating experiments [24, 25]. The morphological evolution presented by the authors is not fully convincing, however. The possibility of displacive austenite formation in the welded joints of low carbon steel was also discussed by RaÈsaÈnen and Tenkula [26]. Based on the observation of particular recrystallization behaviour of heat-aected zones, they suggested the heating rate limit of 3008C/s for the displacive mechanism to be operative. Another interesting aspect of the subject is the formation of austenite from supersaturated solid solutions during continuous heat cycles. It was shown that in microstructures consisting initially of martensite or bainite, austenite forms either at prior austenite grain boundaries or between the laths of martensite/bainitic ferrite where it produces an equiaxed or lath-like/acicular morphology, respectively [27]. The occurrence of one or other morphological form is believed to be due to the degree of superheating and is of practical importance: martensite formed after high-rate short-time austenitizing was found to be much harder than that usually obtained, which was referred to as superhardness [28]. On the other hand, rapidly austenitized bainitic steels can exhibit a so-called ``memory eect'' which hinders the grain re®nement by heat treatment [29]. The reasons for these dierent austenitizing behaviours are not fully understood, but they are likely to have something to do with the mechanism of the transformation. Some early studies suggested that the transformation from martensite to austenite is displacive even at relatively low heating rates [30]. Nevertheless, according to contradictory opinions [8], the mechanism of austenite formation from martensite in carbon steels should be diusion-controlled, because of the rapid precipitation of ®ne carbide particles that are likely to hinder any transformation by shear. The origin of any discrepancy in the discussed area is undoubtedly the variety and complexity of morphologies in the investigated system, and also the number of variables to be considered when analysing phase transformations upon continuous heating. Thus, for better understanding of the austenitization process it should be useful to deal with direct experimental proof based on simple morphological appearances. The aim of this work is to study the morphological aspect of the early stages of austenite formation in a carbon steel during continuous rapid heating. In order to reduce the possibility of interpretation errors due to carbide precipitation or recrystallization phenomena upon heating, the initial structure of annealed ferrite with coarse spheroidized cementite was selected.
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION 2. EXPERIMENTAL PROCEDURE
Strips of carbon steel containing 0.69% C, 0.72% Mn, 0.24% Cr and 0.24% Si were used. In the asdelivered state, the microstructure consisted of ferrite and globular cementite. The shape and size of the specimens were chosen in order to obtain a uniform temperature in the material during heating and cooling. Specimens 100 mm long with a width/ thickness of 15 mm/1 mm were given a preliminary soaking of 30 min at 7008C in order to remove the eects of previous cold work and to obtain equilibrium partition of elements between carbide and bulk phases. The heat cycles were monitored using a special computer-controlled device equipped with water cooling and an infra-red temperature measurement system with an operating range between 400 and 12008C. Heating rates between 200 and 20008C/s were obtained by passing a high intensity electrical alternating current through the specimens. Water quenching was performed by means of an electromagnetic valve and was able to start almost as soon as the maximum temperature of the cycle (Tp) was reached. Figure 1 displays a typical time±temperature recording which further de®nes the parameters of a rapid heat treatment. It should be noted that taking into account both heating and quenching periods between A1 and Tp, the total time involved is 100±200 ms for the heating rates used in this study.
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For each heating rate a series of cycles was made with a maximum temperature from 740 to 8208C so as to observe a large interval of transformation progress. The corresponding microstructure was then veri®ed by light microscopy using chemical etching and interferential contrast techniques. The samples with a rather low transformation degree were ®nally selected for transmission electron microscopy observations. The idea was to ®nd some isolated areas containing both parent and product phases at the early stage of transformation so that standard morphological and crystallographic TEM analysis could be done. Special attention was paid to grain boundary regions and to carbide±ferrite interfaces since these are preferential zones for austenite nucleation. In order to characterize the eect of heating rates on austenitization, conventional water quenching treatments with a heating rate close to 58C/s were also carried out using an electrical furnace. 3. EXPERIMENTAL RESULTS
Optical micrographs illustrating the whole transformation progress are presented in Fig. 2. The irregular areas of dierent size represent here the transformation product against the background of non-transformed ferrite and globular cementite. As revealed by interferential contrast, some of these zones comprise numerous carbide particles which were not dissolved during short-time austenitiza-
Fig. 1. Time vs temperature recording of a rapid heat cycle. The measuring range of the infra-red probe is between 400 and 12008C.
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KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
Fig. 2. Optical micrograph showing the progression of the austenite transformation (Tp=7858C; heating rate = 12008C/s).
tion. At this level of resolution, it should be noted that the outline of the parent/product interface is found to be relatively smooth after short austenitization as well as after furnace treatment. TEM examinations give more information on the behaviour of the interfaces between the two reacting phases. These observations con®rm the earlier reported preferential sites of austenite nucleation because both cementite/ferrite interfaces and ferrite grain boundaries are always found to be at the origin of austenitic areas. Yet, as illustrated by Figs 3 and 4, the aspect of the interface is signi®cantly dierent for the two methods of austenitization. After furnace treatment, the interfaces of the transformed regions clearly exhibit a continuous, regular curved outline [Fig. 3(a)] enclosing one or more carbide particles. Due to water quenching, these zones have obviously experienced one complete a 4 g 4 a' cycle because they were normally transformed into martensite. On the other hand, short time processes usually give rise to a ragged irregular boundary [Fig. 3(b) and Fig. 4] which simply re¯ects the presence of distinctly orientated, lathlike or plate-like components within transformed zones. Sometimes, in well developed regions, usually located at grain corners both aspects of the interface curiously coexist (Fig. 5), as if allotriomorphic and lath-shaped forms of austenite were superposed on one another. Generally, in such zones the lathlike forms are barely visible, as they are partially or entirely entrapped by a regular interface. It is also interesting to note that the growth on both sides of a single grain boundary of ferrite was frequently observed (Fig. 6). Given the cooling conditions, the morphology of a transformed zone clearly corresponds to the shape of its austenite crystal which is characterized by
incomplete growth into the parent grain of ferrite. It was also found that such zones usually start to grow at grain boundaries, near the intergranular particles of cementite. After water cooling, lathshaped areas were sometimes composed of martensite and they also contained a considerable amount of retained austenite as shown in Fig. 6. Most importantly, well isolated, single crystals of austenite entirely preserved down to room temperature were clearly visible. Figures 7 and 8 clearly show single laths of primary, carbide-free austenite growing inside of ferrite grains. Selected area diraction patterns of both parent and product phases could then be taken in order to construct the corresponding stereographic projection. Based on the obtained data the crystallographic orientation relationship between a and g phases is de®ned as well as the growth direction of austenite in the ferritic matrix. The austenite±ferrite orientation relationship observed in this study is within a few degrees of the Kurdjumov±Sachs one: 101g k111a ; 111g k 011a . The major growth direction of austenitic laths was found to be nearly 101g k111a , which corresponds to the close-packed directions of the g and a lattices and to the slip directions. Moreover, carbon concentration of a single austenite lath could be estimated, given the variation of Ms temperature with composition [6, 30±32]. This relationship, con®rmed in several works, has been obtained in real conditions of martensite formation, i.e. it takes into account the actual stress level during the transformation. For a crystal entirely preserved in austenitic state, the corresponding Ms temperature should be the same or lower than the temperature of cooling water. Given the composition (Mn, Cr, Si) of the bulk phase (parent ferrite before transformation to
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
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Fig. 3. General aspect of transformed zone. (a) Furnace treatment: bright ®eld image. (b) Short-time austenitization: Tp=7858C, heating rate = 12008C/s, bright ®eld image. (c) Short-time austenitization: Tp=7858C, heating rate = 12008C/s, (220)g dark ®eld image.
austenite), a value between 0.9 and 1.3% C is found depending on the applied data. This surprisingly high carbon content needs of course further exper-
imental veri®cations. Nevertheless, it seems to indicate an important deviation from equilibrium conditions at this stage of growth. Obviously, car-
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KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
bon concentration of the zones partially or completely transformed to martensite should be lower, probably due to dierent local conditions of carbon diusion.
4. DISCUSSION
Fig. 4. Mounting of TEM micrographs showing numerous austenite laths (partially transformed into martensite) growing on a grain boundary: Tp=7728C, heating rate = 3008C/s, (200)a dark ®eld image.
Fig. 5. Complex aspect of transformed zone (bright ®eld image). Superposition of both allotriomorphic and lathshaped forms of austenite (Tp=7858C, heating rate = 12008C/s).
The present study shows that the morphology of austenite crystals at an early stage of high-rate austenitization is quite similar to that produced during austenite decomposition upon cooling. However, the general conditions of the phase change for both a t g transformations are very dierent according to the nature of the parent and product phases, i.e. a 4 g or g 4 a. Firstly, fast austenitization is carried out at relatively high temperatures compared with those normally reported for martensitic (Ms) and bainitic transformation (Bs), or WidmanstaÈtten ferrite formation (Ws) [6, 31]. Secondly, in view of the strong increase of diusivity with temperature, there are favourable conditions for a reconstructive transformation to occur, unless the driving force created by superheating is high enough to generate the transformation by a displacive mode. On the other hand, at higher temperatures both austenite and ferrite have relatively low yield strength, so plastic accommodation can eciently reduce the strain energy barrier which would impose an a 4 g displacive transformation. In order to propose a general model of austenite formation during rapid heating, it is useful to compare certain principal characteristics of the a 4 g and g 4 a transformation products, including morphological appearance, chemical composition and crystallography of the phases. In this study, for the given range of heating rates we were able to obtain a lath-like transformation product at least in the early stage of austenite growth (Fig. 4). The lath boundaries are always easy to distinguish in the outer parts, whereas they are less pronounced or disappear at the base of the sheaves. Therefore, it seems likely that the longitudinal growth of a sheaf is followed by the coalescence of individual laths giving an impression of their increase in thickness. The transformed zones presented here in Fig. 3(c) and Fig. 4 normally consist of sheaves of laths growing directly from ferrite grain boundaries at a rate of about 70 mm/s. According to the available classi®cations [6, 13] such a morphological form is not equivalent to WidmanstaÈtten ferrite since it would require the development of an allotriomorphic zone prior to secondary sideplates apparition. Furthermore, the similarity with acicular ferrite is also not justi®ed because of dierent nucleation conditions of this morphological modi®cation. It is interesting to mention that WidmanstaÈtten austenite formation has already been reported in low carbon steels [16, 33].
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
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Fig. 6. TEM micrographs indicating growth on both sides of a ferrite grain boundary (Tp=7858C, heating rate = 12008C/s): (a) bright ®eld image; (b) (220)g dark ®eld image; (c) corresponding diraction and key patterns. MÐmartensite, AÐaustenite, FÐferrite.
It is generally admitted that the presence of shear is the main feature of a displacive transformation. In the case of a martensitic one, it brings about a distinct surface relief which is considered as the macroscopic image of the invariant strain. Such a relief was also reported for bainitic transformation and even for WidmanstaÈtten ferrite formation, but its interpretations are very dierent [6, 34±38]. Considering the relatively small volume of lath-shaped austenite in our experiments, it was rather dicult to use the surface relief method in order to con®rm the presence of strain. Instead, TEM observation of mechanical interaction between austenite crystals and cementite particles can give some indication as to the actual strain level that seems to accompany the growth of the laths. As can be seen in Fig. 4, the longitudinal growth of a large sheaf is stopped abruptly
by a coarse cementite particle inside the parent ferrite grain. As a result of this interaction, the central part of the sheaf is considerably deformed at the obstacle which could explain the advanced coalescence of this area. Other examples of the interaction are presented in Fig. 3(c) and Fig. 8. Carbide particles situated directly on the growth axes have came into contact with the lath tips. Such an ``impingement'' gives rise to visible shear of the particles in the vicinity of points of impact, followed by in situ transformation to austenite. It is interesting to mention that in situ austenite formation in cementite has already been predicted theoretically [39] and proved experimentally in steels after laser treatment [24]. Most theoretical considerations of austenite decomposition accept temperature To as the basic parameter for estimating the value of the driving force
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KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
in order to account for the displacive character of the g 4 a reaction [40]. As shown earlier, the initial structure of the material employed in this study consists of ferrite and coarse carbide particles, so we actually examine a single grain of bulk a phase being partially transformed into austenite. The initial chemical composition of the bulk phase is thus automatically de®ned by carbon solubility in ferrite at room temperature, and by the partition coecient of substitutional Mn and Cr atoms between ferrrite and cementite. In order for the transformation to occur by a martensitic mode, the actual To value of the parent phase should be lower than the temperature of the reaction. Moreover, the chemical composition of the transformation product should be the same as that of the parent phase, including
both interstitial and substitutional element contents. According to available data [6], taking account of the initial chemical composition of the ferrite, To of about 8508C is obtained. It is thus higher than the peak temperatures in our experiments and even higher than Ac1 for the given ferrite composition. On the other hand, carbon content in primary laths of austenite seems to dier considerably from the initial composition of the parent ferrite. Therefore, there are two signi®cant reasons for not adopting a martensitic model of transformation. Firstly, the eective driving force is too small for a direct displacive a 4 g transition to take place. Secondly, carbon diusion seems to be somehow involved at the early stage of austenite formation, at least during the nucleation stage.
Fig. 7. Retained austenite formed at a ferrite grain boundary (Tp=7938C, heating rate = 13508C/s): (a) 131g dark ®eld image; (b) corresponding diraction and key patterns. AÐaustenite, FÐferrite.
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
Fig. 8. Retained austenite formed on a cementite particle (Tp=7938C, heating rate = 13508C/s): (a) 111g dark ®eld image; (b) corresponding diraction pattern.
It can be deduced from these observations that a general model closer to bainitic transformation could be accepted, whatever its mechanism (reconstructive or displacive). By comparison with upper bainitic ferrite formation, the following features tend to justify such a model in the case of high-rate austenitization: the carbide-free transformation product has a form of laths that grow directly from grain boundaries and interfaces, showing the KS crystallographic orientation relationship with respect to parent ferrite. The growth rate is similar to that reported for bainitic ferrite formation [29]. The longitudinal growth of the austenite laths is followed by their side-by-side coalescence. Carbon diffusion, generally accepted to take place during bainitic transformation, is also found to play a signi®cant role in the growth of lath-shaped ``bainitic'' austenite upon continuous rapid heating. However, in order to de®ne the exact role of carbon diusion at both nucleation and growth stages, more kinetics data are necessary.
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Clearly, the eective value of the driving force in this case is larger than that required for allotriomorph formation, so the larger superheating due to an increase of heating rate is indispensable. It can be supposed, moreover, that if the driving force is high enough, the transformation could start in a displacive manner before the nucleation process of a reconstructive transformation is completed. In order to ensure favourable thermodynamic conditions, a substantial drop of To should be necessary prior to the growth stage. This situation can arise due to an increase of carbon concentration at grain boundaries, in the presence of coarse carbide particles and due to an easy carbon diusion path. Such favourable thermodynamic conditions can also be easily produced at a/cementite interfaces. Since the volume of the product phase at this stage of reaction is very small, carbon supply to the reaction zone can be ensured without visible dissolution of cementite. According to the well-known models of the interface [41, 42], a grain boundary of a ®nite thickness can be considered as a carbon storage tank in the conditions of a diusion short circuit, by analogy with the collector plate concept [15]. That means for a non-steady state of the growth, the carbon concentration can increase substantially in the interior of a boundary and possibly, in its close vicinity. It is to be noted that the analytic, simpli®ed solution of the problem is only known for the steady state [43] and the theoretical treatment of the present case is rather unrealistic. As mentioned earlier, an increase of heating rate changes the position of Ac1 and Ac3 temperatures. In terms of continuous heating transformation this simply means that the a 4 g reaction is shifted to higher temperatures, although the incubation time is simultaneously reduced. In terms of the Fe±C diagram, this indicates a new position of the Ac1 line that takes the actual heating rate into account. If such a ``dynamic'' Ac1 is plotted on the Fe±C diagram, the a phase ®eld becomes larger and the line of carbon solubility limit can be extrapolated into the two-phase region (Fig. 9). Similar extrapolation is successfully applied to the lines Am, A3 and the solubility of carbon in ferrite when analysing the decomposition of austenite upon cooling. The high heating rate situation is shown schematically in Fig. 9. An Fe±C alloy, initially composed of ferrite with coarse spheroidized cementite, is heated to a temperature Tp with three dierent heating rates. These rates were chosen so as to show three stages of transformation advancement at Tp. When Tp is reached, the transformation is ®nished at the slower rate V1, whereas it only begins at the intermediate one V2 and it has not started at all at the higher rate V3. If V3 is considered, the alloy is continuously heated in the ferritic state up to temperature T3 during a time corresponding at most to the incubation period It. At the same time the solubility limit increases up to the dynamic value (Ld), so some
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KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
Fig. 9. Schematic representation of the eect of heating rate on equilibrium conditions at early stage of austenitization.
supersaturation of the ferritic phase becomes possible. Such a supersaturation is able to produce the driving force necessary to initiate the movement of a ragged, probably semi-coherent interface. However, as previously shown, the resulting morphology of austenite can only be observed at an early stage of growth and the phenomenon is dicult to chart over the whole progress of transformation. Attempts to obtain much of the lath-like phase by an increase of peak temperature were not successful. Instead, the zones that experienced the a 4 g 4 a cycle were completely transformed to very ®ne martensite and in consequence the primary structure of g was no longer visible. This gives the impression that the transformation ceases quickly to advance in the bainitic way and that the further completion of the reaction occurs by migration of a less mobile, probably incoherent interface, at conditions closer to those of local equilibrium. This observation seems to be in close relation with the kinetics data [15, 16] mentioned earlier in this text. The temperature region of 740±7808C given by the authors, corresponds indeed to the peak temperatures that make it possible to obtain the lath-shaped austenite in our study. Otherwise, as shown by TEM observations, individual needles of primary austenite were entirely preserved down to room temperature. It is well known that an enrichment with g-stabilizing elements together with compressive stress are the factors responsible for this behaviour. Taking account of the contraction due to the formation of austenite as well as the dierence of the coecients of thermal expansion of ferrite and austenite, compressive stress cannot be produced at this stage of the growth. On the other hand, if the austenite stabilization is attributed to elevated carbon concentration, the origin of the driving force of austenite formation during continuous rapid heating can be explained. It is dicult to say whether the actual value of the driving force is high enough for a dis-
placive transformation to start. Furthermore, it is beyond the scope of this work to attempt to de®ne the exact thermodynamic conditions of nucleation and growth. Nevertheless, some qualitative characteristics of the growth mechanism of short-time austenitization can be approached. Figure 10 presents a schematic illustration of the possible growth development in three successive steps: as a result of superheating, the partial dissol-
Fig. 10. Schematic illustration of the bainitic austenite formation: (I) grain boundary diusionÐlocal supersaturation; (II) formation of austenite laths; (III) coalescence of austenite lathsЮnal form of the transformed zone.
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
ution of intergranular cementite causes the carbon concentration to increase in a grain boundary zone by easy diusion path (I). Austenite nucleation is thus accompanied by a relatively high driving force, able to induce a rapid movement of a coherent or semi-coherent a$g interface. Consequently, carbon rich austenite laths start to grow at ferrite grain boundaries or a$cementite interfaces (II). At this stage of the reaction, the austenitic phase can be preserved down to room temperature if water quenching is done. At a later step (III) the austenite laths growing in parallel coalesce. Subsequent continuous heating or soaking favours growth by regular interface movement, which progressively causes the primary structure of lath-shaped austenite to vanish. 5. CONCLUSIONS
1. This work shows that a particular lath-shaped morphology is formed at a very early stage of austenitization. The time involved in this phenomenon is shorter than the incubation period of classical reconstructive a 4 g transformation. 2. Under the present conditions of rapid heating, the morphological appearance of austenite at an early stage of the growth is similar to that of upper, carbide-free, bainitic ferrite. 3. The term bainitic austenite is strongly suggested by the following concomitant microstructural features: . the lath-like aspect of the product phase at its early stage of formation; . the direct formation at ferrite grain boundaries; . the Kurdjumov±Sachs orientation relationship with the parent grain of ferrite; . the lateral growth of the sheaves of austenite by side-by-side coalescence. 4. It is postulated that the metastable conditions of growth agree with the extrapolation of the line of carbon solubility in ferrite into the two-phase a + g region of the Fe±C equilibrium diagram. AcknowledgementsÐThe authors would like to thank Jean-Pierre Marteel (Director LAGI2E) for support in realization of the rapid heating device. The technical aid of J. Gest, M. Lareal and F. Marteel is also appreciated. REFERENCES 1. Aaronson, H. I., The Decomposition of Austenite by Diusional Processes, ed. V. F. Zackay and H. I. Aaronson. Interscience, New York, 1962, p. 387. 2. Porter, D. A. and Esterling, K. E., Phase Transformations in Metals and Alloys. Van Nostrand Reinhold, New York, 1981. 3. Dube, C. A., Aaronson, H. I. and Mehl, R. F., Revue MeÂtall., 1958, 55, 201. 4. Spanos, G., Metall. Trans., 1994, 25A, 1967.
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5. Honeycombe, R. W. K. and Bhadeshia, H. K. D. H., Steels, Microstructure and Properties. Arnold, London, 1995. 6. Ohmori, Y., Ohtsubo, H., Jung, Y., Okaguchi, S. and Ohtani, H., Metall. Trans., 1994, 25A, 1981. 7. Krauss, G. and Marder, A. R., Metall. Trans., 1971, 2, 2343. 8. Krauss, G., Steels. ASM Int., Materials Park, 1990. 9. Heheman, E. P., Kinsman, K. R. and Aaronson, H. I., Metall. Trans., 1972, 3, 1077. 10. Hillert, M., Metall. Trans., 1994, 25A, 1957. 11. Aaronson, H. I. and Hall, M. G., Metall. Trans., 1994, 25A, 1797. 12. Christian, J. W. and Edmonds, D. W., Int. Conf. Phase Transformations in Ferrous Alloys, ed. A. R. Marder and J. I. Goldstein. ASM Cleveland, Ohio, 1984, p. 293. 13. Bhadeshia, H. K. D. H. and Christian, J. W., Metall. Trans., 1990, 21A, 767. 14. Hillert, M., Nilson, K. and ToÈrndahl, L. E., J. Iron Steel Inst., 1971, 209, 49. 15. Speich, G. R., Demarest, V. A. and Miller, R. L., Metall. Trans., 1981, 12A, 1419. 16. Jayswal, S. K. and Gupta, S. P., Z. Metallk., 1992, 11, 809. 17. Albutt, K. J. and Garber, S., J. Iron Steel Inst., 1966, 204, 1217. 18. Nikishov, N. A., Sokolov, A. M. and Ulyanov, V. G., Metallov. term. Obrab. Metall., 1991, 8, 2. 19. Morozov, O. P., Physics Metals Metallogr., 1994, 77, 283. 20. Orlich, J., Traitement Thermique, 1974, 90, 69. 21. Schlicht, H., Traitement Thermique, 1974, 89, 63. 22. Zerwekh, R. P. and Wayman, C. M., Acta metall., 1965, 13, 99. 23. Quantin, C., Metauer, G. and Gantois, M., Revue MeÂtall., 1979, 136, 478. 24. Lyasotskii, I. V. and Shtanskii, D. V., Physics Metals Metallogr., 1993, 75, 77. 25. Yakovleva, I. L., Schastlivtsev, V. M., Tabatchikova, T. I., Mirzaiev, D. A. and Osintseva, A. L., Physics Metals Metallogr., 1995, 79, 576. 26. RaÈsaÈnen, E. and Tenkula, J., Scand. J. Metallurgy, 1972, 1, 75. 27. Judd, R. R. and Paxton, H. W., Trans. TMS-AIME, 1968, 242, 206. 28. Scoonover, T. M. and Krauss, G., Met. Eng. Q., 1972, 12, 41. 29. Bhadeshia, H. K. D. H., Bainite in Steels. Institute of Materials, London, 1992. 30. Gulaev, A. P., Metallovedenie. Metallurgia, Moscow, 1977, p. 240. 31. Nishiyama, Z., Martensitic Transformation. Academic Press, London, 1978. 32. Andrews, J. W., J. Iron Steel Inst., 1965, 203, 721. 33. Yi, J. J., Kim, I. S. and Choi, H. S., Metall. Trans., 1985, 16A, 1237. 34. Ko, T. and Cottrell, S. A., J. Iron Steel Inst., 1952, 172, 307. 35. Kinsman, K. R., Eichen, E. and Aaronson, H. I., Metall. Trans., 1975, 6A, 303. 36. Hall, M. G. and Aaronson, H. I., Metall. Trans., 1994, 25A, 1923. 37. Bhadeshia, H. K. D. H., Acta metall., 1981, 29, 1117. 38. Christian, I. W., Metall. Trans., 1990, 21A, 799. 39. Zhukov, A. A., Geometricheskaya Termodynamika Splavov Zheleza. Metallurgia, Moscow, 1979, p. 67. 40. Zener, C., Trans. Am. Inst. Min. Engrs, 1946, 167, 550. 41. Cahn, J. W., Acta metall., 1960, 8, 554. 42. Hillert, M., Metall. Trans., 1975, 6A, 5. 43. Wever, F. and Lange, H., Mitt. Kaiser-Wilhelm Inst. Eisenf., 1932, 14, 71.
Acta mater. Vol. 46, No. 16, pp. 5917±5927, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00210-9 1359-6454/98 $19.00 + 0.00
THE BAINITIC MECHANISM OF AUSTENITE FORMATION DURING RAPID HEATING W. J. KALUBA1{, R. TAILLARD2 and J. FOCT2 Laboratoire d'Applications en GeÂnie Industriel-Energie-Environnement, Universite du Littoral CoÃte d'Opal, 62968 Longuenesse, France and 2Laboratoire de MeÂtallurgie Physique, URA CNRS n8234, Universite de Lille 1, 59655 Villeneuve d'Ascq, France
1
(Received 19 March 1998; accepted 7 May 1998) AbstractÐStrips of a spheroidized 0.68% C, 0.67% Mn and 0.24% Si pearlitic steel were submitted to rapid heat cycles. The samples were heated by direct electrical conduction up to a peak temperature and immediately water quenched. The morphological features of the a 4 g transformation were studied by light and electron microscopy. It was found that austenite at an early stage of the transformation has a speci®c lath-like morphology and that its growth always starts directly from a grain boundary of ferrite. Moreover, the presence of retained austenite allows the prediction of relatively high carbon content in the transformation product. The particular interaction of austenite laths with carbide particles gives an indication that the growth mechanism can be displacive at its early stage. Comparing the morphological characteristics of austenite formation observed in this study with those reported for austenite decomposition, a bainitic model of the early stage of a 4 g transformation is proposed. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. ReÂsumeÂÐLes premiers stades d'austeÂnitisation d'un acier perlitique globularise aÁ 0,68% C, 0,67% Mn et 0,24% Si sont eÂtudieÂs par chauage rapide. De ®nes bandes de meÂtal sont chaueÂes par eet Joule jusqu'aÁ une tempeÂrature maximale avant d'eÃtre immeÂdiatement trempeÂes aÁ l'eau. La morphologie du produit de la transformation a 4 g est eÂtudieÂe par microscopie optique et par microscopie eÂlectronique en transmission. En deÂbut de reÂaction, l'austeÂnite cro|Ãt directement aÁ partir d'un joint de grain ferritique et adopte une morphologie en lattes. La conservation de la nature austeÂnitique de ces lattes aÁ tempeÂrature ambiante suggeÁre son fort enrichissement en carbone. Les interactions observeÂes entre des lattes d'austeÂnite et des particules de carbure militent par ailleurs en faveur d'un meÂcanisme displacif en tout deÂbut de la transformation. Les eÂvolutions morphologiques observeÂes en cours de chauage dans cette eÂtude conduisent aÁ proposer un meÂcanisme de type bainitique pour la transformation a 4 g. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.
1. INTRODUCTION
In solids, the behaviour of the interface between parent and product phases is of chief importance in order to characterize any phase transformation. This behaviour is generally determined by the value of the driving force induced by thermal treatment, and also by the structure of the interface. For instance, an increasing driving force can cause a substantial change of transformation mechanism (from reconstructive to displacive), which usually goes with an evolution of the morphology of the product phase (from equiaxed to lath-like/plate-like in this case). One of the best illustrations of such a tendency is aorded by the widely studied iron±carbon system. The numerous papers on austenite decomposition discuss the relationships between the morphology of ferrite, the mechanism of the g 4 a transformation and the supercooling. The rich variety of resulting morphologies are classi®ed in the following well-known sequence: ferrite allotrio{To whom all correspondence should be addressed.
morphs±WidmanstaÈtten ferrite±bainitic ferrite and martensite at an increasing degree of supercooling [1±8]. Two major transformation mechanisms are reported to take part in the formation of one or other morphological state: a reconstructive one involving a time dependent, diffusional growth and a displacive one based on the movement of coordinated atoms occurring by shear. The reconstructive mechanism is generally accepted to control the growth of ferrite allotriomorphs at low supercooling, whereas the formation of martensite is proved to occur under high supercooling by the displacive manner. However, in spite of the number of works in this area, there are still controversial opinions on the growth mechanism of transformations that takes place at intermediate supercooling, especially concerning the formation of bainite [9±13]. The morphology and mechanism of reverse transformation from ferrite to austenite have been less widely studied, partly because of the number of experimental diculties encountered. The high temperature product phase is dicult to observe
5917
5918
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
directly without the use of sophisticated experimental techniques. Moreover, the morphological features of the transformation are usually modi®ed by cooling to room temperature. Consequently, the problems of austenite formation in the iron±carbon system are often discussed in terms of a well-known isothermal diusion-controlled growth, based on the local equilibrium concept and taking account of the dissolution of cementite in the matrix [14]. The kinetics of austenite formation during intercritical annealing was presented in several works [15, 16]. It was shown that the experimentally measured growth kinetics diers signi®cantly from the theoretical predictions for the temperature range of 740± 7808C and especially for short times of austenitization. It should be noted that the heating method used in these studies (in lead pots or salt bath) as well as the relatively small dimensions of the samples, enables us to predict a relatively high heating rate of the material prior to austenitization. Yet, the in¯uence of the heating rate on the a 4 g transformation at its early stage was not discussed. Moreover, the short times of austenitization (less than 5 s) were not taken into consideration. However, under conditions of continuous heating, especially at high rates, there must be some deviation from the equilibrium model. In close parallel with the well-known eects of cooling, the corresponding increase of the driving force should give rise to a change of the transformation mechanism and in consequence to a morphological evolution of austenite with heating parameters (peak temperature, heating rate . . . ). The problem was the subject of several studies [17±21], including the role of heating rate and the initial microstructure of Fe± C alloys. It was shown that an increasing heating rate brings about a continuous rise of Ac1 and Ac3 temperatures. Moreover, a distinct temperature plateau was observed at a heating rate exceeding some critical value. The occurrence of this plateau was sometimes considered as an indication of a change of the mechanism of transformation from reconstructive to displacive. Unfortunately, the amount of morphological evidence to support such a statement is rather limited, especially for steels with high carbon content. Zerwekh and Wayman [22], using pure iron whiskers, were the ®rst to prove that the a 4 g transformation can occur by a displacive mode even at low overheating. They observed a well-marked shape change during the transformation and de®ned the habit plane. They were not able, however, to observe the same feature on the polycrystalline material. With iron±carbon alloys dierent austenitization behaviours were noticed depending on their initial microstructure. It was suggested that the transformation starts at cementite±ferrite interfaces in a reconstructive way, whereas massive or displacive modes are able to operate at ferrite grain
boundaries [17, 23]. Displacive formation of austenite was later con®rmed in laser heating experiments [24, 25]. The morphological evolution presented by the authors is not fully convincing, however. The possibility of displacive austenite formation in the welded joints of low carbon steel was also discussed by RaÈsaÈnen and Tenkula [26]. Based on the observation of particular recrystallization behaviour of heat-aected zones, they suggested the heating rate limit of 3008C/s for the displacive mechanism to be operative. Another interesting aspect of the subject is the formation of austenite from supersaturated solid solutions during continuous heat cycles. It was shown that in microstructures consisting initially of martensite or bainite, austenite forms either at prior austenite grain boundaries or between the laths of martensite/bainitic ferrite where it produces an equiaxed or lath-like/acicular morphology, respectively [27]. The occurrence of one or other morphological form is believed to be due to the degree of superheating and is of practical importance: martensite formed after high-rate short-time austenitizing was found to be much harder than that usually obtained, which was referred to as superhardness [28]. On the other hand, rapidly austenitized bainitic steels can exhibit a so-called ``memory eect'' which hinders the grain re®nement by heat treatment [29]. The reasons for these dierent austenitizing behaviours are not fully understood, but they are likely to have something to do with the mechanism of the transformation. Some early studies suggested that the transformation from martensite to austenite is displacive even at relatively low heating rates [30]. Nevertheless, according to contradictory opinions [8], the mechanism of austenite formation from martensite in carbon steels should be diusion-controlled, because of the rapid precipitation of ®ne carbide particles that are likely to hinder any transformation by shear. The origin of any discrepancy in the discussed area is undoubtedly the variety and complexity of morphologies in the investigated system, and also the number of variables to be considered when analysing phase transformations upon continuous heating. Thus, for better understanding of the austenitization process it should be useful to deal with direct experimental proof based on simple morphological appearances. The aim of this work is to study the morphological aspect of the early stages of austenite formation in a carbon steel during continuous rapid heating. In order to reduce the possibility of interpretation errors due to carbide precipitation or recrystallization phenomena upon heating, the initial structure of annealed ferrite with coarse spheroidized cementite was selected.
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION 2. EXPERIMENTAL PROCEDURE
Strips of carbon steel containing 0.69% C, 0.72% Mn, 0.24% Cr and 0.24% Si were used. In the asdelivered state, the microstructure consisted of ferrite and globular cementite. The shape and size of the specimens were chosen in order to obtain a uniform temperature in the material during heating and cooling. Specimens 100 mm long with a width/ thickness of 15 mm/1 mm were given a preliminary soaking of 30 min at 7008C in order to remove the eects of previous cold work and to obtain equilibrium partition of elements between carbide and bulk phases. The heat cycles were monitored using a special computer-controlled device equipped with water cooling and an infra-red temperature measurement system with an operating range between 400 and 12008C. Heating rates between 200 and 20008C/s were obtained by passing a high intensity electrical alternating current through the specimens. Water quenching was performed by means of an electromagnetic valve and was able to start almost as soon as the maximum temperature of the cycle (Tp) was reached. Figure 1 displays a typical time±temperature recording which further de®nes the parameters of a rapid heat treatment. It should be noted that taking into account both heating and quenching periods between A1 and Tp, the total time involved is 100±200 ms for the heating rates used in this study.
5919
For each heating rate a series of cycles was made with a maximum temperature from 740 to 8208C so as to observe a large interval of transformation progress. The corresponding microstructure was then veri®ed by light microscopy using chemical etching and interferential contrast techniques. The samples with a rather low transformation degree were ®nally selected for transmission electron microscopy observations. The idea was to ®nd some isolated areas containing both parent and product phases at the early stage of transformation so that standard morphological and crystallographic TEM analysis could be done. Special attention was paid to grain boundary regions and to carbide±ferrite interfaces since these are preferential zones for austenite nucleation. In order to characterize the eect of heating rates on austenitization, conventional water quenching treatments with a heating rate close to 58C/s were also carried out using an electrical furnace. 3. EXPERIMENTAL RESULTS
Optical micrographs illustrating the whole transformation progress are presented in Fig. 2. The irregular areas of dierent size represent here the transformation product against the background of non-transformed ferrite and globular cementite. As revealed by interferential contrast, some of these zones comprise numerous carbide particles which were not dissolved during short-time austenitiza-
Fig. 1. Time vs temperature recording of a rapid heat cycle. The measuring range of the infra-red probe is between 400 and 12008C.
5920
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
Fig. 2. Optical micrograph showing the progression of the austenite transformation (Tp=7858C; heating rate = 12008C/s).
tion. At this level of resolution, it should be noted that the outline of the parent/product interface is found to be relatively smooth after short austenitization as well as after furnace treatment. TEM examinations give more information on the behaviour of the interfaces between the two reacting phases. These observations con®rm the earlier reported preferential sites of austenite nucleation because both cementite/ferrite interfaces and ferrite grain boundaries are always found to be at the origin of austenitic areas. Yet, as illustrated by Figs 3 and 4, the aspect of the interface is signi®cantly dierent for the two methods of austenitization. After furnace treatment, the interfaces of the transformed regions clearly exhibit a continuous, regular curved outline [Fig. 3(a)] enclosing one or more carbide particles. Due to water quenching, these zones have obviously experienced one complete a 4 g 4 a' cycle because they were normally transformed into martensite. On the other hand, short time processes usually give rise to a ragged irregular boundary [Fig. 3(b) and Fig. 4] which simply re¯ects the presence of distinctly orientated, lathlike or plate-like components within transformed zones. Sometimes, in well developed regions, usually located at grain corners both aspects of the interface curiously coexist (Fig. 5), as if allotriomorphic and lath-shaped forms of austenite were superposed on one another. Generally, in such zones the lathlike forms are barely visible, as they are partially or entirely entrapped by a regular interface. It is also interesting to note that the growth on both sides of a single grain boundary of ferrite was frequently observed (Fig. 6). Given the cooling conditions, the morphology of a transformed zone clearly corresponds to the shape of its austenite crystal which is characterized by
incomplete growth into the parent grain of ferrite. It was also found that such zones usually start to grow at grain boundaries, near the intergranular particles of cementite. After water cooling, lathshaped areas were sometimes composed of martensite and they also contained a considerable amount of retained austenite as shown in Fig. 6. Most importantly, well isolated, single crystals of austenite entirely preserved down to room temperature were clearly visible. Figures 7 and 8 clearly show single laths of primary, carbide-free austenite growing inside of ferrite grains. Selected area diraction patterns of both parent and product phases could then be taken in order to construct the corresponding stereographic projection. Based on the obtained data the crystallographic orientation relationship between a and g phases is de®ned as well as the growth direction of austenite in the ferritic matrix. The austenite±ferrite orientation relationship observed in this study is within a few degrees of the Kurdjumov±Sachs one: 101g k111a ; 111g k 011a . The major growth direction of austenitic laths was found to be nearly 101g k111a , which corresponds to the close-packed directions of the g and a lattices and to the slip directions. Moreover, carbon concentration of a single austenite lath could be estimated, given the variation of Ms temperature with composition [6, 30±32]. This relationship, con®rmed in several works, has been obtained in real conditions of martensite formation, i.e. it takes into account the actual stress level during the transformation. For a crystal entirely preserved in austenitic state, the corresponding Ms temperature should be the same or lower than the temperature of cooling water. Given the composition (Mn, Cr, Si) of the bulk phase (parent ferrite before transformation to
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
5921
Fig. 3. General aspect of transformed zone. (a) Furnace treatment: bright ®eld image. (b) Short-time austenitization: Tp=7858C, heating rate = 12008C/s, bright ®eld image. (c) Short-time austenitization: Tp=7858C, heating rate = 12008C/s, (220)g dark ®eld image.
austenite), a value between 0.9 and 1.3% C is found depending on the applied data. This surprisingly high carbon content needs of course further exper-
imental veri®cations. Nevertheless, it seems to indicate an important deviation from equilibrium conditions at this stage of growth. Obviously, car-
5922
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
bon concentration of the zones partially or completely transformed to martensite should be lower, probably due to dierent local conditions of carbon diusion.
4. DISCUSSION
Fig. 4. Mounting of TEM micrographs showing numerous austenite laths (partially transformed into martensite) growing on a grain boundary: Tp=7728C, heating rate = 3008C/s, (200)a dark ®eld image.
Fig. 5. Complex aspect of transformed zone (bright ®eld image). Superposition of both allotriomorphic and lathshaped forms of austenite (Tp=7858C, heating rate = 12008C/s).
The present study shows that the morphology of austenite crystals at an early stage of high-rate austenitization is quite similar to that produced during austenite decomposition upon cooling. However, the general conditions of the phase change for both a t g transformations are very dierent according to the nature of the parent and product phases, i.e. a 4 g or g 4 a. Firstly, fast austenitization is carried out at relatively high temperatures compared with those normally reported for martensitic (Ms) and bainitic transformation (Bs), or WidmanstaÈtten ferrite formation (Ws) [6, 31]. Secondly, in view of the strong increase of diusivity with temperature, there are favourable conditions for a reconstructive transformation to occur, unless the driving force created by superheating is high enough to generate the transformation by a displacive mode. On the other hand, at higher temperatures both austenite and ferrite have relatively low yield strength, so plastic accommodation can eciently reduce the strain energy barrier which would impose an a 4 g displacive transformation. In order to propose a general model of austenite formation during rapid heating, it is useful to compare certain principal characteristics of the a 4 g and g 4 a transformation products, including morphological appearance, chemical composition and crystallography of the phases. In this study, for the given range of heating rates we were able to obtain a lath-like transformation product at least in the early stage of austenite growth (Fig. 4). The lath boundaries are always easy to distinguish in the outer parts, whereas they are less pronounced or disappear at the base of the sheaves. Therefore, it seems likely that the longitudinal growth of a sheaf is followed by the coalescence of individual laths giving an impression of their increase in thickness. The transformed zones presented here in Fig. 3(c) and Fig. 4 normally consist of sheaves of laths growing directly from ferrite grain boundaries at a rate of about 70 mm/s. According to the available classi®cations [6, 13] such a morphological form is not equivalent to WidmanstaÈtten ferrite since it would require the development of an allotriomorphic zone prior to secondary sideplates apparition. Furthermore, the similarity with acicular ferrite is also not justi®ed because of dierent nucleation conditions of this morphological modi®cation. It is interesting to mention that WidmanstaÈtten austenite formation has already been reported in low carbon steels [16, 33].
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
5923
Fig. 6. TEM micrographs indicating growth on both sides of a ferrite grain boundary (Tp=7858C, heating rate = 12008C/s): (a) bright ®eld image; (b) (220)g dark ®eld image; (c) corresponding diraction and key patterns. MÐmartensite, AÐaustenite, FÐferrite.
It is generally admitted that the presence of shear is the main feature of a displacive transformation. In the case of a martensitic one, it brings about a distinct surface relief which is considered as the macroscopic image of the invariant strain. Such a relief was also reported for bainitic transformation and even for WidmanstaÈtten ferrite formation, but its interpretations are very dierent [6, 34±38]. Considering the relatively small volume of lath-shaped austenite in our experiments, it was rather dicult to use the surface relief method in order to con®rm the presence of strain. Instead, TEM observation of mechanical interaction between austenite crystals and cementite particles can give some indication as to the actual strain level that seems to accompany the growth of the laths. As can be seen in Fig. 4, the longitudinal growth of a large sheaf is stopped abruptly
by a coarse cementite particle inside the parent ferrite grain. As a result of this interaction, the central part of the sheaf is considerably deformed at the obstacle which could explain the advanced coalescence of this area. Other examples of the interaction are presented in Fig. 3(c) and Fig. 8. Carbide particles situated directly on the growth axes have came into contact with the lath tips. Such an ``impingement'' gives rise to visible shear of the particles in the vicinity of points of impact, followed by in situ transformation to austenite. It is interesting to mention that in situ austenite formation in cementite has already been predicted theoretically [39] and proved experimentally in steels after laser treatment [24]. Most theoretical considerations of austenite decomposition accept temperature To as the basic parameter for estimating the value of the driving force
5924
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
in order to account for the displacive character of the g 4 a reaction [40]. As shown earlier, the initial structure of the material employed in this study consists of ferrite and coarse carbide particles, so we actually examine a single grain of bulk a phase being partially transformed into austenite. The initial chemical composition of the bulk phase is thus automatically de®ned by carbon solubility in ferrite at room temperature, and by the partition coecient of substitutional Mn and Cr atoms between ferrrite and cementite. In order for the transformation to occur by a martensitic mode, the actual To value of the parent phase should be lower than the temperature of the reaction. Moreover, the chemical composition of the transformation product should be the same as that of the parent phase, including
both interstitial and substitutional element contents. According to available data [6], taking account of the initial chemical composition of the ferrite, To of about 8508C is obtained. It is thus higher than the peak temperatures in our experiments and even higher than Ac1 for the given ferrite composition. On the other hand, carbon content in primary laths of austenite seems to dier considerably from the initial composition of the parent ferrite. Therefore, there are two signi®cant reasons for not adopting a martensitic model of transformation. Firstly, the eective driving force is too small for a direct displacive a 4 g transition to take place. Secondly, carbon diusion seems to be somehow involved at the early stage of austenite formation, at least during the nucleation stage.
Fig. 7. Retained austenite formed at a ferrite grain boundary (Tp=7938C, heating rate = 13508C/s): (a) 131g dark ®eld image; (b) corresponding diraction and key patterns. AÐaustenite, FÐferrite.
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
Fig. 8. Retained austenite formed on a cementite particle (Tp=7938C, heating rate = 13508C/s): (a) 111g dark ®eld image; (b) corresponding diraction pattern.
It can be deduced from these observations that a general model closer to bainitic transformation could be accepted, whatever its mechanism (reconstructive or displacive). By comparison with upper bainitic ferrite formation, the following features tend to justify such a model in the case of high-rate austenitization: the carbide-free transformation product has a form of laths that grow directly from grain boundaries and interfaces, showing the KS crystallographic orientation relationship with respect to parent ferrite. The growth rate is similar to that reported for bainitic ferrite formation [29]. The longitudinal growth of the austenite laths is followed by their side-by-side coalescence. Carbon diffusion, generally accepted to take place during bainitic transformation, is also found to play a signi®cant role in the growth of lath-shaped ``bainitic'' austenite upon continuous rapid heating. However, in order to de®ne the exact role of carbon diusion at both nucleation and growth stages, more kinetics data are necessary.
5925
Clearly, the eective value of the driving force in this case is larger than that required for allotriomorph formation, so the larger superheating due to an increase of heating rate is indispensable. It can be supposed, moreover, that if the driving force is high enough, the transformation could start in a displacive manner before the nucleation process of a reconstructive transformation is completed. In order to ensure favourable thermodynamic conditions, a substantial drop of To should be necessary prior to the growth stage. This situation can arise due to an increase of carbon concentration at grain boundaries, in the presence of coarse carbide particles and due to an easy carbon diusion path. Such favourable thermodynamic conditions can also be easily produced at a/cementite interfaces. Since the volume of the product phase at this stage of reaction is very small, carbon supply to the reaction zone can be ensured without visible dissolution of cementite. According to the well-known models of the interface [41, 42], a grain boundary of a ®nite thickness can be considered as a carbon storage tank in the conditions of a diusion short circuit, by analogy with the collector plate concept [15]. That means for a non-steady state of the growth, the carbon concentration can increase substantially in the interior of a boundary and possibly, in its close vicinity. It is to be noted that the analytic, simpli®ed solution of the problem is only known for the steady state [43] and the theoretical treatment of the present case is rather unrealistic. As mentioned earlier, an increase of heating rate changes the position of Ac1 and Ac3 temperatures. In terms of continuous heating transformation this simply means that the a 4 g reaction is shifted to higher temperatures, although the incubation time is simultaneously reduced. In terms of the Fe±C diagram, this indicates a new position of the Ac1 line that takes the actual heating rate into account. If such a ``dynamic'' Ac1 is plotted on the Fe±C diagram, the a phase ®eld becomes larger and the line of carbon solubility limit can be extrapolated into the two-phase region (Fig. 9). Similar extrapolation is successfully applied to the lines Am, A3 and the solubility of carbon in ferrite when analysing the decomposition of austenite upon cooling. The high heating rate situation is shown schematically in Fig. 9. An Fe±C alloy, initially composed of ferrite with coarse spheroidized cementite, is heated to a temperature Tp with three dierent heating rates. These rates were chosen so as to show three stages of transformation advancement at Tp. When Tp is reached, the transformation is ®nished at the slower rate V1, whereas it only begins at the intermediate one V2 and it has not started at all at the higher rate V3. If V3 is considered, the alloy is continuously heated in the ferritic state up to temperature T3 during a time corresponding at most to the incubation period It. At the same time the solubility limit increases up to the dynamic value (Ld), so some
5926
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
Fig. 9. Schematic representation of the eect of heating rate on equilibrium conditions at early stage of austenitization.
supersaturation of the ferritic phase becomes possible. Such a supersaturation is able to produce the driving force necessary to initiate the movement of a ragged, probably semi-coherent interface. However, as previously shown, the resulting morphology of austenite can only be observed at an early stage of growth and the phenomenon is dicult to chart over the whole progress of transformation. Attempts to obtain much of the lath-like phase by an increase of peak temperature were not successful. Instead, the zones that experienced the a 4 g 4 a cycle were completely transformed to very ®ne martensite and in consequence the primary structure of g was no longer visible. This gives the impression that the transformation ceases quickly to advance in the bainitic way and that the further completion of the reaction occurs by migration of a less mobile, probably incoherent interface, at conditions closer to those of local equilibrium. This observation seems to be in close relation with the kinetics data [15, 16] mentioned earlier in this text. The temperature region of 740±7808C given by the authors, corresponds indeed to the peak temperatures that make it possible to obtain the lath-shaped austenite in our study. Otherwise, as shown by TEM observations, individual needles of primary austenite were entirely preserved down to room temperature. It is well known that an enrichment with g-stabilizing elements together with compressive stress are the factors responsible for this behaviour. Taking account of the contraction due to the formation of austenite as well as the dierence of the coecients of thermal expansion of ferrite and austenite, compressive stress cannot be produced at this stage of the growth. On the other hand, if the austenite stabilization is attributed to elevated carbon concentration, the origin of the driving force of austenite formation during continuous rapid heating can be explained. It is dicult to say whether the actual value of the driving force is high enough for a dis-
placive transformation to start. Furthermore, it is beyond the scope of this work to attempt to de®ne the exact thermodynamic conditions of nucleation and growth. Nevertheless, some qualitative characteristics of the growth mechanism of short-time austenitization can be approached. Figure 10 presents a schematic illustration of the possible growth development in three successive steps: as a result of superheating, the partial dissol-
Fig. 10. Schematic illustration of the bainitic austenite formation: (I) grain boundary diusionÐlocal supersaturation; (II) formation of austenite laths; (III) coalescence of austenite lathsЮnal form of the transformed zone.
KALUBA: BAINITIC MECHANISM OF AUSTENITE FORMATION
ution of intergranular cementite causes the carbon concentration to increase in a grain boundary zone by easy diusion path (I). Austenite nucleation is thus accompanied by a relatively high driving force, able to induce a rapid movement of a coherent or semi-coherent a$g interface. Consequently, carbon rich austenite laths start to grow at ferrite grain boundaries or a$cementite interfaces (II). At this stage of the reaction, the austenitic phase can be preserved down to room temperature if water quenching is done. At a later step (III) the austenite laths growing in parallel coalesce. Subsequent continuous heating or soaking favours growth by regular interface movement, which progressively causes the primary structure of lath-shaped austenite to vanish. 5. CONCLUSIONS
1. This work shows that a particular lath-shaped morphology is formed at a very early stage of austenitization. The time involved in this phenomenon is shorter than the incubation period of classical reconstructive a 4 g transformation. 2. Under the present conditions of rapid heating, the morphological appearance of austenite at an early stage of the growth is similar to that of upper, carbide-free, bainitic ferrite. 3. The term bainitic austenite is strongly suggested by the following concomitant microstructural features: . the lath-like aspect of the product phase at its early stage of formation; . the direct formation at ferrite grain boundaries; . the Kurdjumov±Sachs orientation relationship with the parent grain of ferrite; . the lateral growth of the sheaves of austenite by side-by-side coalescence. 4. It is postulated that the metastable conditions of growth agree with the extrapolation of the line of carbon solubility in ferrite into the two-phase a + g region of the Fe±C equilibrium diagram. AcknowledgementsÐThe authors would like to thank Jean-Pierre Marteel (Director LAGI2E) for support in realization of the rapid heating device. The technical aid of J. Gest, M. Lareal and F. Marteel is also appreciated. REFERENCES 1. Aaronson, H. I., The Decomposition of Austenite by Diusional Processes, ed. V. F. Zackay and H. I. Aaronson. Interscience, New York, 1962, p. 387. 2. Porter, D. A. and Esterling, K. E., Phase Transformations in Metals and Alloys. Van Nostrand Reinhold, New York, 1981. 3. Dube, C. A., Aaronson, H. I. and Mehl, R. F., Revue MeÂtall., 1958, 55, 201. 4. Spanos, G., Metall. Trans., 1994, 25A, 1967.
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