Observations concerning transformation interfaces in steels

Acta metall, mater. Vol. 41, No. 12, pp. 3421-3434, 1993 0956-7151/93 $6.00 + 0.00 Printed in Great Britain. All rights reserved Copyright © 1993 Pergamon Press Ltd

OBSERVATIONS CONCERNING TRANSFORMATION INTERFACES IN STEELS

F. A. KHALID and D. V. E D M O N D S t

Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, England

(Received I0 January 1993; in revised form 26 April 1993)

Akaraet--The transformation interfaces of pearlite, allotilomorphic cementite, M23C6, and Widmanst/itten cementite plates in high-Mn high-C alloy steels have been studied by TEM. Linear striations in the interface have been analysed and related to intersections with stacking faults in the parent austenite phase. Emphasis is given to the pearlite interface where it is found that the striations at the interface increased as a result of thermomechanical treatment of the austenite prior to isothermal transformation, consistent with an increased density of planar defects. The effect of heat treatment, and Si alloying additions, are also considered. Both conventional and in situ TEM of the pearlite interface showed that the linear defects stretched across both ferrite and cementite phases at the pearlite interface, apparently without any deviation or change in image contrast. The results are compared with similar ones made of the static ),/~ interphase boundaries in duplex stainless steel. The effect of prior deformation structure in the parent austenite on the growth and interface structure of Widmanst/itten cementite plates has also been considered.

Rrsumr----Nous avons 6tudi6 en microscopic 61ectronique h transmission les interfaces de transformation de la perlite, de la ckmentite allotriomorphe, M23 C6, et de prrcipitrs de ckmentite de Widmanst~itten, dans des aciers fi forte teneur en manganrse et en carbone. Les stiles linraires des interfaces ont 6t6 analysres et relires aux drfauts d'empilement prrsents dans la phase parente. Nous nous sommes intrressrs plus particulirrement aux interfaces de la perlite; nous avons constat6 que la densit6 de stiles fi l'interface augmente si raustrnite a subi un traitement thermomrcanique prralablement ~i la transformation thermique, ce qui est cohrrent avec une densit6 accrue de drfauts plans. Nous avons 6galement examin6 reffet de traitements thermiques et de l'alliage avec du silicium. La microscopic conventionnelle aussi bien qu'in situ montrent que les drfauts linraires se prolongent a la fois dans la ferilte et la ckmentite, sans drviation apparente, ni modification du contraste. Nous avons compar6 ces r~sultats avec des observations analogues d'interfaces ~/], dans des aciers biphasrs inoxydables. Nous avons 6galement 6tudi6 l'effet de structures de drformation prrexistantes dans I'austrnite parente sur la croissance des prrcipitrs de Widmanst/itten.

ZusmnmenfassRg--Die Reaktionsfronten des Perlits, allotropen Zementits, M23C6 und der Widmannst/itten-Zementitplatten in Mn- und C-reichen St~hlen wurden mit Hilfe der Durchstrahlung- selektronen-mikroskopie (TEM) untersucht. Lineare Riefen, die durch das Schneiden der Grenzfl/ichen mit Stapelfehlern im Austenit hervorgingen, wurden analysiert. Besondere Bedeutung wird der Perlitgrensfl/iche zugeschileben. Hier zeigt sich, daft die Riefen in der Grenzfliiche infolge einer thermo- mechanischen Vorbehandlung des Austenits vor der isothermen Umwandlung zunahmen. Dies stimmt fiberein mit der erhrhten planaren Defektdichte. Die Auswirkungen der W/irmebehandlung und das Zulegieren von Si wurden ebenfalls betrachtet. Sowohl konventionelle als auch in-situ TEM-Unter- suchungen der Peditgrenzfl~icbe zeigen, daft sich die linearen Defekte fiber die Ferrit- und Zementitphasen an der Perlitgrenztt/iche erstrecken--anscheinend ohne Abweichung oder ,~nderung des Bildkontrastes. Die Ergebnisse werden mit /ihnlichen Untersuchungen von statischen 7/ct Phasengrenzen an rostfreien Duplex-St/ihlen verglichen. Weiterhin wurde der Einflu8 der thermomechaniscben Vorbehandlung auf das Wachstum der Widmannst/itten-Zementitplatten untersucht.

INTRODUCTION

The structure of t ransformation interfaces, particu- larly those in steels, is normally impossible to study because the interface is destroyed during cooling from the reaction temperature to room temperature. In steels this is usually due to martensitic transformation of the remaining untransformed austenite regions. This problem has been partly circumvented

tPresent address: School of Materials, University of Leeds, Leeds LS2 9JT, England

by using high-Mn high-C steel that remains austenitic at room temperature, but which decomposes pro- gressively during annealing at elevated temperatures. The remaining untransformed austenite and the t ransformation interface are retained at room temperature, thus allowing their examination by transmission electron microscopy.

This alloy system has been used to study morphology and phase orientation relationships between austenite and proeutectoid cementite [1-9] and pearlite [10]. Recent studies carried out on the austenite/pearlite interface have examined the interface

3421

3422 KHALID and EDMONDS: TRANSFORMATION INTERFACES IN STEELS

structure and observed sets of linear striations or defects in the interface. These have been interpreted by Hackney and Shiflet [11-14] and Zhou and Shiflet [15] as growth steps or ledges, and their lateral movement in the interface plane has been proposed as a mechanism for the edgewise growth of pearlite. It has also been suggested, on the basis of TEM examination, that the ledges move across both ferrite

and cementite phases at the pearlite interface, despite the fact that these two phases have quite different crystal structures. Hackney and Shiflet [13] propose that the ability of the pearlite growth ledges to affect the position of the three phase triple junction allows the coupling of the thermodynamic and/or kinetic parameters controlling the development of lamellar pearlite with the atomic scale growth mechanism.

Fig. 1. (a) Coarsely-spaced dislocations interacting with the matrix dislocations BF, (b) DFr, (c) a fine set of intrinsic dislocations (arrows) WBDF g = [200] and (d) stacking faults in the austenite interacting

with the ),/~t interphase boundary, DF = g[200]r, steel 1.

KHALID and EDMONDS: TRANSFORMATION INTERFACES IN STEELS 3423

The presence of two or more different sets of mobile ledges at the growth interface allows the quick and efficient alteration of the growth direction, plate thickness and spacing through the initiation and propagation of ferrite/cementite direction steps. Zhou and Shiflet [15] also examined the steps at the austenite/pearlite interface and the interlamellar interfaces using lattice imaging techniques. They explained that the interlamellar steps are associated with curvature and result directly from pearlite growth ledges. These interpretations are radically different to the previous generally accepted explanations for the growth of pearlite [16, 17].

This experimental alloy steel has also been used to determine phase orientation relationships between austenite and Widmanstatten cementite plates by Pitsch [2] and more recently new orientations have been determined by Thompson and Howell [18, 19], Farooque and Edmonds [9] and Zhou and Shiflet [20]. The transformation characteristics of proeutectoid Widmanstatten cementite have also been studied by Cowley and Edmonds [6] and Khalid et al. [21] and suggested to be not inconsistent with a displacive transformation mechanism. They also observed linear defects which exhibited periodic image contrast along the austenite/cementite interface. Recently, Spanos and Aaronson [7, 8] have examined the interfacial structure and also confirmed the habit plane of proeutectoid cementite plates in a similar alloy steel.

In this paper the interfaces of the three main transformation products; pearlite, grain boundary allotriomorphs and intragranular Widmanstatten cementite, in high-Mn alloy steels are examined. The features at the transformation interfaces of these austenite decomposition products in high-alloy steels are compared with the defect structure observed at the static 7/a boundaries of duplex stainless steel. The pearlite transformation interface has been examined using both conventional and in situ hot-stage TEM. The growth interface of pearlite in deformed austenite has also been examined in order to study the role of planar defects in austenite ahead of the interface. The structure and morphology of grain b o u n d a r y M 2 3 C 6 carbides also observed in these high-Mn steels is also considered.

E X P E R I M E N T A L

The chemical compositions of the alloy steels used for this investigation are given in Table 1. Duplex stainless steel 1 was used in the as hot-rolled con- dition. Alloy steel 2 was made as a 50 g argon arc

melt using high purity materials. The ingot was homogenized at 1250°C for 80 h after encapsulating in silica tube under a partial pressure of argon gas. The ingot was then hot rolled, swaged and machined to 3 mm diameter rod. Alloy steel 3 was made as a 20 kg vacuum melt, forged and then hot rolled to 10 mm diameter. 3 mm diameter rod was machined from the rolled bar. Isothermal heat treatment was carried out on discs of 1 mm thickness cut from the 3 mm diameter rod. Specimens were heat treated in evacuated silica tube under a partial pressure of argon gas. Specimens were solution treated either at 1200 or 1000°C for 60min and quenched in water, solutionised again at 1200 or 1000°C for 15 min and isothermally transformed between 630 and 460°C for 10-18h and quenched in water. Austenitised and water quenched samples of the 3 mm diameter rod were also deformed in tension (to 8% reduction) at room temperature in order to generate a high density of defects in austenite prior to isothermal transformation to pearlite and Widmanstatten cementite. Thin foils for TEM examination were prepared from disc specimens by mechanical grinding down to 90/~m thickness and then electropolishing using a twin-jet Fischione electropolishing unit and a solution of 100 g chro- mium trioxide in 500 ml of glacial acetic acid at an applied potential of 100V. The thin foil specimens were examined in Philips CM 12 and CM 20 electron microscopes operated at 120 and 200 kV, respectively. In situ hot stage TEM experiments were per- formed in a JEOL 2000FX electron microscope operated at 200 kV.

R E S U L T S A N D D I S C U S S I O N

I. Characterization o f the 7/ct interphase boundary in Duplex stainless steel

Figure l(a) and (b) show a general austenite/ferrite microstructure observed in steel 1. Matrix dislocations can be observed interacting with the static ~,/~ grain boundary to give coarsely-spaced defects within the boundary. Fine unidirectional arrays of grain boundary dislocations (arrows) were also observed at the ~/~t boundaries [Fig. l(c)], with spacings of the order of 2-10 nm. Similar features have been observed previously [22-24]. A new observation is that some of the coarse-spaced linear defects are connected with stacking faults present in the austenite as shown in Fig. l(d). The nature of similar non- periodic arrays at the transformation interfaces of pearlite, grain boundary carbides and Widmanstatten cementite plates is considered in more detail in the following section.

Table 1. Chemical compositions of alloy steels

Elements (wt%)

C Mn Si Cr Ni Mo Cu N

Steel 1 0.025 1.41 0.32 22.53 5.94 2.75 0.18 0.152 Steel 2 0.81 11.70 . . . . . . Steel 3 0.96 12.85 0.78 . . . . .


(d) show a small number of irregularly spaced long parallel lines of contrast at the austenite/ferrite interface, and also at the austenite/cementite interface of the pearlite [Fig. 3(e)]. Evidence of defects can also be noticed at the ferrite/cementite product interface [Fig. 3(d)]; however, these did not appear to be connected with the straight defects at the growth interface. The linear striations can also be observed to continue, apparently without distortion, across both ferrite and cementite phases at the interface, e.g. Fig. 3(e). These linear features at the interface appear to be similar to those previously reported by Hackney and Shiflet [11-13].

Figure 4 shows another example of defects visible at the austenite/ferrite interface and dislocations present in the austenite ahead of the interface. There is some evidence for connectivity between these

Fig. 2. (a) Austenite/pearlite interface in alloy 1, BF and corresponding convergent beam electron diffraction CBED patterns. A--austenite, F--ferrite and C----cementite, steel 2.

2. OBSERVATIONS ON THE STRUCTURE OF TRANSFORMATION INTERFACES

2.1. Pearlite

2.1.1. Austenite /pearlite interface. Figure 2 shows the stable austenite/pearlite transformation interface in alloy steel 2 after isothermal transformation, along with convergent beam electron diffraction (CBED) patterns of the respective phases, and no evidence of martensite, whereas, in the case of commercial steels, the interface would normally be destroyed by the formation of martensite during cooling to room temperature. SAD pattern analysis of the pearlite nodules nucleated on the austenite grain boundary reveals that the pearlitic ferrite has a Kurdjumov- Sachs (K-S) or Nishiyama-Wasserman (N-W) orientation relationship [25-27] to that austenite grain into which it is not growing, and no simple orientation relationship with the austenite grain into which it is growing. This is consistent with the results of previous studies [I0].

Weak beam TEM [28] examination has revealed the presence of a finely spaced set of intrinsic dislocations at the austenite/ferrite interface [Fig. 3(a,b)]. A change in the image contrast is visible at the interaction with planar faults and a variation in the spacings of the dislocations with the curvature of the interface is also evident (arrows). Figure 3(c) and Fig. 3(a) and (b). Caption on facing page.


Fig. 3. (a) Examples of austenite/pearlite interface showing (a) a set of finely spaced intrinsic dislocations, BF, (b) WBDF g = [110L, and (c) lines of strong contrast at the austenite/ferrite and ferrite/cementite interfaces, BF, (d) WBDF, g =[011L inset SAD pattern, and (e) linear defects across both the

austenite/ferrite and austenite/cementite interfaces (arrows), BF.

dislocations in the parent austenlte, and the defects imaged in the transformation interface. Conse- quently, there is some similarity with the extrinsic boundary dislocations imaged in the static grain boundaries of duplex stainless steel or in other single phase systems [29, 30].

However, the majority of deformation structure visible in the parent austenite is in the form of dissociated dislocations. These have been analysed as intrinsic stacking faults lying on { 111 } austenite planes (Fig. 5). If the stacking faults in the austenite and the linear striations in the transformation inter-

face are imaged simultaneously, then it can be shown convincingly that the interfacial features arise from intersection of the interface with these planar defects in the austenite.

The possibility exists that the features described above are created after formation of the pearlite, for example, by the heat treatment procedures, or even by the specimen preparation methods. Consequently, further experiments were conducted in order to investigate the origins of the defects.

2.1.2. Influence of heat treatment. In specimens austenitised at a lower temperature, 1000°C as

A M 41,12- -G


content, TEM examination also revealed a higher density of planar defects intersected by the interface, even after heat treatment at the higher temperature of 1200°C. Figure 8 shows both the high density of linear defects at the interface and the larger concentration of stacking faults intersecting the interface. Figure 8(c) shows an example of two sets of defects at the transformation interface.

2.1.5. In situ hot stage TEM. In this experiment austenite/pearlite interfaces of alloy steel 2 were studied, to examine the relationship between the austenite defects and the linear striations at the interface. Thin foil specimens were heated at 470°C for 50 s, cooled to 25°C, and the image recorded. Repeating this procedure a number of times, revealed no significant change in the austenite defect concentration or the image contrast at the interface after each heating period [Fig. 9(b,c)]. However, when the same interface was heated at 500°C for 90 s it was observed that the interface had advanced to give an increment of pearlite growth, and that this was

Fig. 4. (a) Austenite/pearlite interface showing another example of linear defects and (b) WBDF g = [01T]~.

compared to 1200°C, an increased density of dislocations and stacking faults in the parent austenite was obtained. This can be attributed to a slower rate of annihilation of stacking faults at the lower austenitis- ing temperature. A corresponding increase in the number of linear features observed at the austenite/pearlite interface was found. This provides evidence that this defect structure exists, at least in part, in the untransformed parent austenite. In addition, specimens cooled at different rates from the solution treatment/isothermal ageing temperatures, for example, water quenched, air cooled, and furnace cooled, appeared to show little change in the appar- ent density of defects either in the parent austenite phase or intersecting with the interface, and so it can further be concluded that the defects are present at the temperature of formation of the pearlite.

2.1.3. Influence of deformed austenite on the interface. The effect of prior deformation of austenite, to produce a higher density of planar defects in the parent phase before partial transformation to pearlite, has been examined. Figure 6(a) shows nucleation of pearlite nodules at austenite grain boundaries and also at twin boundaries after prior deformation at room temperature. Figure 6(b) shows the increased concentration of planar defects in austenite intersecting with both the ferrite and cementite interfaces of pearlite. Figure 7 shows another example of the increased number of linear defects at the interface as a result of the high density of planar defects produced in austenite by the deformation.

2.1.4. Composition effect. In alloy steel 3, contain- ing a Si addition which lowers the stacking fault energy in austenite, and with a slightly higher C

Fig. 5. Pearlite interface intersecting with the intrinsic stacking faults present in the parent austenite (a) BF, (b) DF

g = [200]~, (c) SAD pattern from the austenite.


irregular fringes where the change in spacing may be associated with curvature of the interface [Fig. 10(c)]. It appears that the interfacial structure of these allotriomorphic carbide phases in steel are similar to allotriomorphic ferrite and other interphase interfaces both in steels and certain non-ferrous alloys [24, 31,32].

Fig. 6. (a) Nucleation of pearlite at the austenite grain boundaries and twin boundaries and (b) planar defects in the austenite intersecting with the austenite/ferrite and

austenite/cementite interfaces, BF.

apparently accompanied by an increase in the defect concentration in the austenite ahead of the interface and the appearance of new linear striations in the interface [Fig. 9(d)]. It would thus appear that the defect concentration in the austenite ahead of the interface is partly created by the transformation stresses due to the volume change involved. The interface then clearly has to advance through these defects resulting in the intersections observed.

2.2. Grain boundary carbides

2.2.1. Allotriomorphic cementite interface. The interface structure of allotriomorphic cementite nucleating on the austenite grain boundaries has also been examined. The allotriomorphs nucleate with a Pitsch relationship [5] to one austenite grain and grow into the adjacent grain with which they have no orientation relationship [10]. Figure 10 shows dislocations and stacking faults in the austenite grain into which the cementite allotriomorph is growing, with evidence of the stacking faults intersecting the cementite interface to produce linear striations. However, this interface also shows other features that could be associated either with interface structure or image contrast effects. High-resolution TEM examination of this interface reveals two sets of fringes at the interface; one set of straight fringes with uniform spacing in the range of 2-4 nm, and a second set of

Fig. 7. Austenite/pearlite interface showing an increase in the linear defects due to high density of the planar defects produced by deformation in austenite, (a) BF, (b), DF, g = [011]~ and (c) interface tilted to show intersecting planar

defects in the austenite, BF.


2.2.2. Me3C6 sawteeth carbide. A sawteeth morphology of carbide at the grain boundaries was observed in alloy steel 3 as illustrated in Fig. 11. Diffraction analysis revealed this to be f.c.c. M23C6 carbide (M = Fe, Mn) and not orthorhombic cementite [Fig. l l(b and c)]; these carbides displayed a cube--cube orientation relationship with respect to one of the austenite grains (A2) into which they grow. These observations are similar to recently reported work on high Mn steels [33-35]. The development of facets and macroscopic steps at the growth interface of the grain boundary carbide was also observed (Fig. 12). Planar faults observed in the carbide are identified as stacking faults lying along { l 11 } planes and these are connected with linear features at the planar interfaces.

2.3. Intragranular Widmanstatten cementite plate interface

Intragranular cementite plates were also obtained at lower heat treatment temperatures and were found to obey the well known Pitsch orientation relationship with the austenite matrix [4, 5] as illustrated in Fig. 13 although new orientation relationships have also been reported recently [9, 18-20]. It is also evident that stacking faults in the austenite intersect with the broad faces of the Widmanstatten cementite plates resulting in coarse linear striations in the interface as observed for the other forms of cementite decomposition products (Fig. 14). It was found that these features at the broad faces of the plate are partial dislocations lying on { 111 } planes of austenite. In addition, other defect structures can be observed in the interface: extrinsic dislocations originating from the austenite were also observed. Finer and more regularly spaced sets of fringes can also be detected; Fig. 15 is a weak beam TEM image showing a regularly spaced fine set of fringes indicating an intrinsic dislocation structure (small arrows) at the broad face of the Widmanstatten cementite which have a spacing of the order of 3-8 nm. Evidence for similar interfacial features have also been reported for lath and plate martensite by Sandvik and Wayrnan [36, 37]. More coarsely spaced but regular lines of strong contrast are also visible at the same interface (large arrows). Other features are also observed but may have resulted from intersection with defects present within the body of the plate [Fig. 15(a)]. Spanos and Aaronson [8] have identified similar features at the interface as sets of finely spaced straight ledges and large kinked ledges, but no refer- ence to the planar defects either in the austenite matrix or in the body of cementite plates was men- tioned. In other work studying the precipitation of Cu in similar alloy steels [38], it has been observed that the copper particles can nucleate at coarsely spaced linear defects on the broad faces of the Widmanstatten cementite plates. This suggests that these linear lines of contrast reveal stationary facets rather than growth ledges.

Fig. 8. Austenite/pearlite interface showing intersection of planar defects in austenite in steel 3, (a) B,F, (b) DF, g = [200]~ and (c) two sets of defects observed at the pearlite

interface (arrows), BF.

The effect of a high density of defects produced by room temperature mechanical treatment of the austenite, on the formation of Widmanstatten cementite plates has also been examined. Figure 16 shows formation of cementite plates in deformed austenite. A high concentration of defects is found in the body of the plates and the plate morphology is perturbated by its growth through the deformed structure [Fig. 16(a)]. This result is in accordance with previous observations reported for the formation of bainitic plates by Sandvik [39]. A large concentration of stacking fault intersections with the broad faces of the plate is evident. Figure 16(b) is an example revealing two sets of extrinsic defects (arrows A and B) at the broad face of the plate which have resulted from


interaction with a high density of stacking faults in the austenite. This was not generally found at the broad faces of cementite plates in the undeformed specimens. However, this is consistent with the results observed in the case of pearlite interfaces formed in deformed austenite. There is also evidence to indicate that the planar defects in the austenite were trapped in the body of the plate [Fig. 16(c)]. This was not observed in the case of the pearlite transformation, which showed a defect-free transformation product (e.g. see Fig. 7) after growth into deformed austenite. This evidence lends further support to the view that growth of cementite plates is not inconsistent with displacive transformation as reported previously [6, 211.

SUMMARY AND CONCLUSIONS

Detailed TEM examination of the interfaces of the three transformation products of austenite in hyper- eutectoid alloy steels has been carried out. Obser- vations at the pearlite growth interface have shown that coarsely-spaced linear defects were not intrinsic growth ledges on these interfaces, but extrinsic defects resulting from intersection of the growing interface with planar stacking faults present in the parent austenite. When the stacking faults intersect the ferrite and cementite at the interface simultaneously, they can result in combined or continuous defects across both the ferrite and cementite interfaces. Similar evidence for y/~ interphase boundaries

Fig. 9. In situ hot stage TEM micrographs of austenite/pearlite interface, (a) interface prior to heating, (b) after heating at 470°C for 50 s, (c) again heated at 470°C for 50 s, and (d) shows an increment of pearlite growth, and also an increase in the defect concentration in the austenite after heating at 500°C for 80 s,

steel 2.


observed in the duplex stainless steel also confirms the observations of the transformation interface. There is evidence to suggest that the austenite con- tains an array of such defects which is enhanced near the interface due to the accompanying transform-

ation stresses, which is also supported by in situ TEM observations.

Different sets of parallel lines of image contrast were also observed at the interfaces of allotriomorphic grain boundary cementite, and intragranular

Fig. I0. Grain boundary cementite showing intersections with the planar defects present in the austenite and fine fringe contrast visible at the interface, (a) BF, (b) DF, g = [111]~ and (c) high-resolution TEM micrograph showing resolvable (001) cementite planes and linear features at the austenite/grain boundary

cementite interface (inset SAD pattern), steel 2.


Fig. 11. M23C 6 carbide exhibiting a sawteeth morphology and cube-cube orientation relationship with the austenite grain A 2, (a) BF, and (b) and (c) SAD patterns, Z = [011]

and [112] respectively.

Widmanstat ten cementite plates. The structure at the Widmansta t ten plate interface showed five types of line contrast (i) finely spaced regular sets (ii) coarsely- spaced regular sets (iii) extrinsic dislocations from the matrix (iv) lines connected with defects within the body of the plate, and (v) linear striations due to intersection with stacking faults in the matrix. Furthermore, the concentration of linear defects increased at the broad faces of the plates in the case of deformed austenite. Planar defects in the austenite

Fig. 13. (a) Widmanstatten cementite plates displayed a Pitsch orientation relationship with respect to the austenite, BF and (b) SAD pattern and corresponding analysis,

steel 2.

were also found trapped in the body of the cementite plates during growth, unlike in the case for pearlite.

M23 C6 carbide occurred at the grain boundaries in addition to the allotriomorphic cementite formation.

Fig. 12. Macroscopic steps and fine facets at the austenite/grain boundary carbide interface interacting with stacking faults in the carbide, BF.


Fig. 14. Examples showing interface structure of the plate (a) intersection of stacking faults at the broad faces and (b) planar stacking faults in both austenite and within the body of the plate, BF.

Evidence of stacking faults in the austenite intersecting with the interfaces of these carbides was also observed.

It is clear that the use of these high-Mn steels can reveal information on the structure of the transformation interfaces in steels. A number of linear features with either fine spacings, or coarser more irregular spacings, can be imaged. In the case of proeutectoid cementite, whether as an allotriomorphic grain

boundary phase, pearlite, or as an intragranular Widmanstatten plate, both fine and coarse sets of images are found, that may be identified as intrinsic or extrinsic structural features accordingly, and these are similar to observations made in this study, or reported by others, of the interfacial structures of grain boundaries in single phase austenite steels, or in the 7/ct boundaries of duplex steels.

Fig. 15. Austenite/Widmanstatten cementite interface revealing (i) fine regular sets, (ii) coarsely spaced regular sets, (iii) extrinsic dislocations from matrix, (iv) lines connected with defects within the body of the plate and (v) linear striations due to interactions with stacking faults, (a) BF and (b) WBDF g = [200]¢.


Fig. 16. (a) Perturbations in a cementite plate formed in highly deformed austenite, (b) showing two sets of extrinsic dislocations (arrows A and B) at the broad face of cementite plate DF g = [200]: and

(c) showing interfacial structure and planar defects trapped in the body of the plate, BF.

Acknowledgements--The authors wish to thank Professor Sir Peter Hirsch F.R.S. for the provision of laboratory facilities and Dr M. L. Jenkins and Mr R. C. Doole for arranging hot stage TEM at BP Research Sunbury, where Dr D. White and Mr A. Scott are thanked for their help during the hot stage TEM experiments. F. A. Khalid is grateful to the Ministry of Science and Technology for an S&T scholarship and PCSIR Laboratories Lahore, Pakistan, for sponsorship.

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Observations concerning transformation interfaces in steels