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Comparison of interphase boundary structure between bainite and martensite in steel
Scripta Materialia 47 (2002) 193–199 www.actamat-journals.com
Comparison of interphase boundary structure between bainite and martensite in steel T. Moritani a, N. Miyajima a, T. Furuhara
b,*
, T. Maki
b
a
b
Graduate Student, Kyoto University, Sakyo-ku, Kyoto 606-8501, Japan Department of Materials Science and Engineering, Yoshida-honmachi, Kyoto University, Sakyo-ku, Kyoto 606-8501, Japan
Abstract Interphase boundary structures were compared between lath-shaped bainitic ferrite in Fe–0.6C–2Si–1Mn and lath martensite in Fe–20Ni–5.5Mn. Both laths exhibit the macroscopic habit plane scattered around (1 2 1)c which contains monoatomic steps with the ð1 1 1Þckð0 1 1Þa terrace (transformation dislocations/structural ledges). Similar accommodation dislocations, with pure-screw characters on ð1 1 1Þckð0 1 1Þa, were observed in both cases. Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Phase transformation; Bainite; Martensite; Austenite; Crystallography; Interface
1. Introduction Atomic structure of interphase boundary provides important information concerning nucleation and growth mechanisms in phase transformations. In martensitic (diffusionless/displacive) transformations, martensite/matrix interphase boundary is fully glissile and can move conservatively [1]. For a0 (bcc) lath martensite formed in c(fcc) austenite in ferrous alloys, dislocation slip occurs as lattice invariant deformation, resulting in a glissile arrangement of interfacial dislocations. So far, various glissile c=a0 interphase boundary models were proposed [2–5]. Sandvik and Wayman [6] first studied the lath martensite/austenite boundaries in
* Corresponding author. Tel.: +81-75-753-5469; fax: +81-75753-4861. E-mail address: [email protected] (T. Furuhara).
Fe–20%Ni–5%Mn by means of transmission electron microscopy (TEM). They observed purescrew dislocations, which should move by glide, on the atomic habit plane (i.e., the parallel close packed planes: ð1 1 1Þckð0 1 1Þa0 ). They reported the presence of similar screw dislocations in Fe– 8%Cr–1%C [7]. Later, Mahon et al. [8] showed, with high-resolution transmission electron microscopy (HREM), that the broad face of lath martensite contains transformation dislocations (monoatomic steps with the ð1 1 1Þckð0 1 1Þa0 terrace) in Fe–8%Cr–1%C. Bainite transformation in steels occurs in the intermediate temperature range between diffusional ferrite/pearlite transformations and martensitic transformation and its mechanism is still subject to controversy [9,10]. Aaronson et al. [11] proposed that sessile misfit dislocations should exist on the interphase boundary formed during diffusional transformation. In the previous TEM study on bainite/austenite boundaries [12,13],
1359-6462/02/$ - see front matter Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 2 ) 0 0 1 2 8 - 8
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sessile dislocations were observed on the broad face of bainitic ferrite. However, no HREM study has been performed on the bainite/austenite boundary and the characteristic of accommodation dislocations is not fully understood. The present authors recently studied interphase boundaries of both martensite and bainite by means of HREM [14,15]. In the present study, the characters of accommodation dislocations on those two kinds of interphase boundaries are compared to discuss the mechanism of boundary migration.
2. Experimental procedure Fe–20.2Ni–5.4Mn (mass%), in which austenite is stable at room temperature and partly transforms isothermally to lath martensite below room temperature [16,17], was used to study lath martensite/austenite boundaries. After austenitized at 1473 K for 3.6 ks and water quenched, the specimens were held at 223 K for various periods to promote isothermal martensitic transformation. Fe–2.0Si–1.0Mn–0.59C (mass%) was used to study bainitic ferrite/austenite boundaries. Specimens were austenitized at 1423 K for 0.6 ks, transformed at 723 K for various periods and water quenched. In this alloy, austenite can be obtained in the gap between adjacent parallel bainitic ferrite laths since
carbon is enriched in austenite without cementite precipitation [18]. Microstructures were observed by means of TEM (Philips CM200) and HREM (Jeol JEM4000EX). Orientation relationships (ORs) between austenite and product were examined by analyzing Kikuchi patterns.
3. Results 3.1. Lath martensite/austenite interphase boundary Martensite holds the ORs slightly scattered around K–S ðð1 1 1Þckð0 1 1Þa0 ; ½1 0 1ck½1 1 1a0 Þ, N ðð1 1 1Þckð0 1 1Þa0 ; ½1 1 0ck½1 0 0a0 Þ and G–T relationships with respect to austenite. The habit plane of lath deviates between ð1 1 1Þcðkð0 1 1Þa0 Þ and ð1 2 1Þcðkð1 3 2Þa0 Þ. In the dark-field image of Fig. 1(a), straight dislocations with an average spacing of 4.8 nm are observed on the broad face of lath. The stereographic projection of Fig. 1(b) shows the crystallographic information for this interface. Close packed planes of two phases are misoriented by 1.0° whereas angular deviation of close packed directions is 3.7° for this interface. The Burgers vector of dislocations was determined to be b1 ¼ a=2½0 1 1c ¼ a=2½1 1 1a0 by contrast analysis. Trace analysis revealed that the microscopic line direction of dislocations is ½0:60; 0:57; 0:57a0 on the atomic habit plane ð1 1 1Þckð0 1 1Þa0 , indicating
Fig. 1. (a) Dark-field TEM micrographs of the dislocations (see white lines) on the broad face of lath martensite (223 K, 28.8 ks transformed), (b) ½0 0 1a0 stereographic projection showing the OR, the habit plane and the nature of dislocations for the interface in (a).
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that those dislocations are in a pure-screw orientation within an error of 2°. This result is consistent with the observation made by Sandvik and Wayman [6]. In HREM observation, two different beam directions were chosen; the parallel close-packed directions ½1 0 1ck½1 1 1a0 of the K–S OR and ½1 1 0ck½1 0 0a0 which are parallel in the N OR. Fig. 2 shows the HREM images of the broad face of martensite lath observed along these two directions. The broad face of lath, edge-on along ½101ck½111a0 (Fig. 2(a)), contains regularly spaced monoatomic steps with the ð111Þckð011Þa0 terrace as previously reported [8]. These steps can be regarded as transformation dislocations. Along ½110ck½100a0 (Fig. 2(b)), it is seen that many accommodation dislocations with Burgers vectors lying on the parallel close packed planes are present on the interface. At the edge of lath, such a dislocation existed per every third layer of the parallel close packed planes. Fig. 3(a) schematically shows the shear strain originated from the stacking sequence change of
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parallel close packed planes. A strain of a= 12½1 1 2cð¼ a=6½0 1 1a0 Þ, which is perpendicular to ½1 1 0ck½1 0 0a0 , is associated with the stacking change per one ð1 1 1Þc layer. By the coalescence of six transformation dislocations, the shear strain of a=2½1 1 2c is accumulated. This strain can be accommodated by two kinds of perfect dislocations; b1 ¼ a=2½0 1 1cð¼ a=2½1 1 1a0 Þ, b2 ¼ a=2½1 0 1cð¼ a=2½1 1 1aÞ. By introducing these two dislocations alternatively on every third layer of (1 1 1)c, the shear strain due to the stacking change of Fig. 3(a) is fully accommodated. There is another strain associated with the shape change of parallel close packed planes. To accommodate this component, as shown in Fig. 3(b), the same two sets of dislocations b1 and b2 (the solid lines in the figure) need to be introduced on the ð1 1 1Þckð0 1 1Þa0 plane. Although both of those dislocations are almost of pure-screw type on the interface when N OR is held across the boundary, they can be visualized as extra half planes when they are viewed along ½1 1 0ck½1 0 0a0 . Suppose that those dislocations exist as loops on the slip planes inclined on the parallel close packed planes. They can accommodate both of the strain due to the shape change on the parallel close packed planes (i.e., the atomistic habit plane of the broad face) and the strain arising from stacking sequence change at the edge, simultaneously. The dislocations observed in Fig. 2(b) can be explained well by the combination of those two kinds of dislocations. Thus, it is concluded that the shear strain due to the change in the stacking and the shape of parallel close packed planes are fully accommodated by such an arrangement of interphase boundary dislocations. 3.2. Bainitic ferrite/austenite interphase boundary
Fig. 2. HREM micrographs showing the broad face of lath martensite viewed along (a) ½1 0 1ck½1 1 1a0 and (b) ½1 1 0ck½1 0 0a0 , respectively.
Bainitic ferrite formed at 723 K is lath shaped with a thickness of about 1 lm. Each lath consists of smaller sub-units whose thickness is in the order of 0.1 lm, as previously reported [19,20]. Sub-units of ferrite exhibit the ORs scattered around K–S, N and G–T relationships and its habit plane deviates around ð1 2 1Þcðkð1 3 2Þa0 Þ. In the dark-field TEM micrograph of Fig. 4(a), straight dislocations, 5 nm spaced, are present on
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Fig. 3. (a) Stacking sequence change for the parallel close packed planes viewed along ½1 1 0ck½1 0 0a0 and (b) atomic matching on the ð1 1 1Þckð0 1 1Þa0 plane when N OR is held.
Fig. 4. (a) Dark-field TEM micrograph showing the dislocations on the broad face of bainitic ferrite lath (723 K 250 s transformed), (b) ½0 0 1a stereographic projection representing the OR, the habit plane and the nature of dislocations for the interface in (a).
the broad face of bainitic ferrite lath. The stereographic projection of Fig. 4(b) shows that the close packed planes of the two phases are misoriented by 1.5° whereas the angular deviation between the close packed directions is less than 1°. The Burgers vector of dislocations determined by contrast analysis is b1 ¼ a=2½0 1 1c ¼ a=2½1 1 1a0 , which is the same as martensite. Trace analyses of the habit plane and the line direction of dislocations revealed that these dislocations are in a purescrew orientation on the atomic habit plane ð1 1 1Þckð0 1 1Þa0 within an error of less than 2°. The observed dislocation structure was compared with the arrangement of dislocations with no residual strain predicted by O-lattice calculation. In the present study, the same approach as
the study by Hall et al. [21] was used for appropriate lattice parameter in the material used. Fig. 5 shows the calculated dislocation structures on the conjugate parallel close packed plane for ORs varying from the K–S to the N relationship. In the case of N OR (Fig. 5(a)), both sets of accommodation dislocations, are nearly in screw orientations. On the other hand, for the K–S relationship (Fig. 5(c)), the dislocation with b1 ¼ a=2½0 1 1cð¼ a=2½1 1 1aÞ has a mixed character whereas the dislocation with b2 ¼ a=2½1 0 1cð¼ a=2½1 1 1aÞ remains nearly a screw type. For the OR which corresponds to the boundary shown in Fig. 4, the dislocation with b1 is of a mixed type (Fig. 5(b)). The angular deviation between the Burgers vector and the line direction of dislocation D1 is as much
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Fig. 5. O-lattice calculation for the ð1 1 1Þckð0 1 1Þa interface with a misorientation ðhÞ between ½1 0 1c and ½1 1 1a: (a) h ¼ 5:26° (N OR); (b) h ¼ 1° (the OR for Fig. 4); (c) h ¼ 0° (K–S OR). The Burger vectors are b1 ¼ a=2½0 1 1cð¼ a=2½1 1 1aÞ and b2 ¼ a=2½1 0 1cð¼ a=2½1 1 1aÞ.
as 42°. This line direction is significantly different from the observation, implying that misfit strain is not fully accommodated on the ð1 1 1Þckð0 1 1Þa interface. Fig. 6 shows the atomic structures on the broad face of bainitic ferrite. In Fig. 6(a), the macroscopic habit plane consists of monoatomic steps identical to the transformation dislocations observed on the martensite/austenite boundary shown in Fig. 2(a). In Fig. 6(b), accommodation dislocations similar to those on the boundary of lath martensite in Fig. 2(b) are also visible. It is clear that the bainitic ferrite/austenite boundary structure is similar to those of lath martensite with residual transformation strain.
On the contrary to the observation of Fig. 4, sessile bainitic ferrite/austenite boundary structures were reported previously [12,13] in good agreement with the equilibrium dislocation structure predicted by a geometrical model [22]. Some boundaries containing the dislocations in sessile orientations were also recognized in the present study, of which detail is reported elsewhere [23]. Furthermore, it was also proposed [12] that the broad face of bainitic ferrite contains triatomic structural ledges to accommodate misfit in an elastic manner. In the present observation, however, triatomic or higher steps always accompanied some accommodation dislocations.
4. Discussion
Fig. 6. HREM micrographs showing the broad face of bainitic ferrite viewed along (a) ½1 0 1ck½1 1 1a and (b) ½1 1 0ck½1 0 0a.
In the present study, similar pure-screw accommodation dislocations on ð1 1 1Þckð0 1 1Þa were observed in both cases of lath martensite and bainitic ferrite. One of the present authors [24] proposed that bainitic ferrite has a glissile array of interphase boundary dislocations in the early stage. The pure-screw dislocations presently observed on the broad face of bainitic ferrite appears consistent with that proposal but contradicts to the previous study reporting the presence of sessile dislocations on the broad face of bainitic ferrite. The question now arises why both sessile and glissile arrangements of accommodation dislocations are observed on the bainitic ferrite/austenite boundary. The present authors consider that the nature of interphase boundary depends on degrees of strain
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relaxation. There are two strain components in the atomic structure change across the ð1 1 1Þckð0 1 1Þa boundary. One originates from the shape change as described in Fig. 3(b) and the other is the volumetric strain arising from the difference in atomic density on the plane. When interphase boundary diffusion is involved in the stress relaxation around particles, only the strain due to shape change can be relaxed, resulting in a hydrostatic stress field acting on the boundary [25]. When there is a difference in atomic density between the two phases, long-range diffusion is necessary to relax that component. Thus, the transport of extra atoms by volume diffusion of substitutional atoms is required for boundary migration. It should take extremely longer time than the isothermal holding time employed in the present study for bainitic ferrite to grow to the observed sizes. One may consider that the metastable glissile boundary structure changes to the equilibrium sessile structure after sub-units of bainitic ferrite stop growing by some reasons, e.g., accumulation of elastic strain. Such modification process was previously proposed [26–29] although it was regarded that they seem infeasible [30]. Transition of boundary structure should occur by diffusional stress relaxation. If accommodation was aided by short-circuit diffusion, the relaxation time required could be in the order of isothermal transformation time. Since many intruder dislocations are attached to the boundary due to plastic deformation in bainitic ferrite and austenite, fast diffusion along such dislocations might partly affect to the kinetics of strain accommodation. On the transition of interphase boundary structure, however, there can be another point of view which is completely opposite to the one discussed above. Suppose that a sessile boundary is formed during bainitic transformation. During quenching from transformation temperature, the driving force for transformation increases and finally exceeds the critical one necessary for martensitic growth. Under such a circumstance, the equilibrium sessile boundary structure might change to a metastable glissile one. Thus, further consideration is necessary to clarify the detail of strain accommodation process.
Another and more important question to be addressed is whether the boundary observed with two sets of screw dislocations are truly glissile. Christian and Knowles [31] made the proposal that the glissile boundary containing two sets of dislocations are either with two different Burgers vectors on the same glide plane or with the same Burgers vector but on different glide planes. The combination of two sets of screw dislocations in Fig. 3(b) is sessile when their glide planes are different each other as described in the preceding section. If so, the interphase boundary of lath martensite observed becomes sessile. Sandvick and Wayman [7] proposed the operation of the same glide plane for the two sets of dislocations with different Burgers vectors to avoid conflicts. However, the glide planes of those dislocations still need to be identified in future studies. 5. Summary The present study clearly indicates that similar boundary structures can be formed during martensitic and bainite transformations in steels. Acknowledgements The authors are grateful to Professor Hubert I. Aaronson (Carnegie Mellon University, USA) and Professor Gary R. Purdy (McMaster University, Canada) for stimulating discussions. References [1] Clark HM, Wayman CM. In: Aaronson HI, editor. Phase transformations. Metals Park, Ohio: ASM; 1970. p. 59– 114. [2] Frank FC. Acta Metall 1953;1:15–21. [3] Suzuki H. Sci Rep Tohoku Univ 1954;6:30–50. [4] Olson GB, Cohen M. Metall Trans A 1976;7A:1905–14. [5] Sandvik BPJ, Wayman CM. Metall Trans A 1983;14A: 835–44. [6] Sandvik BPJ, Wayman CM. Metall Trans A 1983;14A: 823–34. [7] Sandvik BPJ, Wayman CM. Metall Trans A 1983;14A: 2455–77. [8] Mahon GJ, Howe JM, Mahajan S. Philos Mag Lett 1989;59:273–9.
T. Moritani et al. / Scripta Materialia 47 (2002) 193–199 [9] Bhadeshia HKDH, Christian JW. Metall Trans A 1990; 21A:767–98. [10] Aaronson HI, Reynolds Jr WT, Shiflet GJ, Spanos G. Metall Trans A 1990;21A:767–98. [11] Aaronson HI, Furuhara T, Rigsbee JM, Reynolds Jr WT, Howe JM. Metall Trans A 1990;21A:2369–409. [12] Rigsbee JM, Aaronson HI. Acta Metall 1979;27:365–76. [13] Li C, Perovic V, Purdy GR. In: Lorimer GW, editor. Proceedings of the Conference on Phase Transformations’87. London: Institute of Metals; 1988. p. 326–9. [14] Miyajima N, Moritani T, Furuhara T, Maki T. In: Proceedings of the 4th Pacific Rim International Conference on Advanced Materials and Processing (PRICM4). The Japan Institute of Metals; 2001. p. 2507–10. [15] Furuhara T, Moritani T, Miyajima N, Maki T. To be presented in the International Conference on Martensitic Transformations (ICOMAT’02), Helsinki, Finland, 2002. [16] Wakasa K, Wayman CM. Acta Metall 1981;29:973–90. [17] Wakasa K, Wayman CM. Acta Metall 1981;29:991– 1011. [18] Tsuzaki K, Nakao C, Maki T. Mater Trans, JIM 1991;32: 658–66. [19] Oblak JM, Hehemann RF. In: Transformation and hardenability in steels. Ann Arbor, MI: Climax Molybdenum Co.; 1967. p. 15–30.
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[20] Tsuzaki K, Kodai A, Maki T. Metall Trans A 1994;25A: 2009–16. [21] Hall MG, Rigsbee JM, Aaronson HI. Acta Metall 1986;34:1419–31. [22] Rigsbee JM, Aaronson HI. Acta Metall 1979;27:351–63. [23] Moritani T, Miyajima N, Furuhara T, Maki T. Mater Trans, in preparation. [24] Tsuzaki K, Maki T. J de Phys IV, Colloque C8 1995: 61–70. [25] Onaka S, Kato M. Mater Sci Eng 1991;A146:217–32. [26] Hehemann RF. In: Aaronson HI, editor. Phase transformations. Metals Park, Ohio: ASM; 1970. p. 397–432. [27] Bhadeshia HKDH, Edmonds DV. Acta Metall 1980;28: 1265–73. [28] Ohmori Y, Ohtsubo H, Jung YC, Okaguchi S, Ohtani H. Metall Mater Trans A 1994;25A:1981–9. [29] Ohmori Y, Jung Y-C, Nakai K, Shioiri H. Acta Mater 2001;49:3149–62. [30] Rigsbee JM, Menon ESK, Lee HJ, Aaronson HI. Scripta Metall 1983;7:1465–73. [31] Christian JW, Knowles KW. In: Aaronson HI, Laughlin DE, Sekerka RF, Wayman CM, editors. Proceedings of the International Conference on Solid-to-Solid Phase Transformations. Warrendale, PA: TMS-AIME; 1982. p. 1185– 207.
Comparison of interphase boundary structure between bainite and martensite in steel T. Moritani a, N. Miyajima a, T. Furuhara
b,*
, T. Maki
b
a
b
Graduate Student, Kyoto University, Sakyo-ku, Kyoto 606-8501, Japan Department of Materials Science and Engineering, Yoshida-honmachi, Kyoto University, Sakyo-ku, Kyoto 606-8501, Japan
Abstract Interphase boundary structures were compared between lath-shaped bainitic ferrite in Fe–0.6C–2Si–1Mn and lath martensite in Fe–20Ni–5.5Mn. Both laths exhibit the macroscopic habit plane scattered around (1 2 1)c which contains monoatomic steps with the ð1 1 1Þckð0 1 1Þa terrace (transformation dislocations/structural ledges). Similar accommodation dislocations, with pure-screw characters on ð1 1 1Þckð0 1 1Þa, were observed in both cases. Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Phase transformation; Bainite; Martensite; Austenite; Crystallography; Interface
1. Introduction Atomic structure of interphase boundary provides important information concerning nucleation and growth mechanisms in phase transformations. In martensitic (diffusionless/displacive) transformations, martensite/matrix interphase boundary is fully glissile and can move conservatively [1]. For a0 (bcc) lath martensite formed in c(fcc) austenite in ferrous alloys, dislocation slip occurs as lattice invariant deformation, resulting in a glissile arrangement of interfacial dislocations. So far, various glissile c=a0 interphase boundary models were proposed [2–5]. Sandvik and Wayman [6] first studied the lath martensite/austenite boundaries in
* Corresponding author. Tel.: +81-75-753-5469; fax: +81-75753-4861. E-mail address: [email protected] (T. Furuhara).
Fe–20%Ni–5%Mn by means of transmission electron microscopy (TEM). They observed purescrew dislocations, which should move by glide, on the atomic habit plane (i.e., the parallel close packed planes: ð1 1 1Þckð0 1 1Þa0 ). They reported the presence of similar screw dislocations in Fe– 8%Cr–1%C [7]. Later, Mahon et al. [8] showed, with high-resolution transmission electron microscopy (HREM), that the broad face of lath martensite contains transformation dislocations (monoatomic steps with the ð1 1 1Þckð0 1 1Þa0 terrace) in Fe–8%Cr–1%C. Bainite transformation in steels occurs in the intermediate temperature range between diffusional ferrite/pearlite transformations and martensitic transformation and its mechanism is still subject to controversy [9,10]. Aaronson et al. [11] proposed that sessile misfit dislocations should exist on the interphase boundary formed during diffusional transformation. In the previous TEM study on bainite/austenite boundaries [12,13],
1359-6462/02/$ - see front matter Ó 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 6 2 ( 0 2 ) 0 0 1 2 8 - 8
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sessile dislocations were observed on the broad face of bainitic ferrite. However, no HREM study has been performed on the bainite/austenite boundary and the characteristic of accommodation dislocations is not fully understood. The present authors recently studied interphase boundaries of both martensite and bainite by means of HREM [14,15]. In the present study, the characters of accommodation dislocations on those two kinds of interphase boundaries are compared to discuss the mechanism of boundary migration.
2. Experimental procedure Fe–20.2Ni–5.4Mn (mass%), in which austenite is stable at room temperature and partly transforms isothermally to lath martensite below room temperature [16,17], was used to study lath martensite/austenite boundaries. After austenitized at 1473 K for 3.6 ks and water quenched, the specimens were held at 223 K for various periods to promote isothermal martensitic transformation. Fe–2.0Si–1.0Mn–0.59C (mass%) was used to study bainitic ferrite/austenite boundaries. Specimens were austenitized at 1423 K for 0.6 ks, transformed at 723 K for various periods and water quenched. In this alloy, austenite can be obtained in the gap between adjacent parallel bainitic ferrite laths since
carbon is enriched in austenite without cementite precipitation [18]. Microstructures were observed by means of TEM (Philips CM200) and HREM (Jeol JEM4000EX). Orientation relationships (ORs) between austenite and product were examined by analyzing Kikuchi patterns.
3. Results 3.1. Lath martensite/austenite interphase boundary Martensite holds the ORs slightly scattered around K–S ðð1 1 1Þckð0 1 1Þa0 ; ½1 0 1ck½1 1 1a0 Þ, N ðð1 1 1Þckð0 1 1Þa0 ; ½1 1 0ck½1 0 0a0 Þ and G–T relationships with respect to austenite. The habit plane of lath deviates between ð1 1 1Þcðkð0 1 1Þa0 Þ and ð1 2 1Þcðkð1 3 2Þa0 Þ. In the dark-field image of Fig. 1(a), straight dislocations with an average spacing of 4.8 nm are observed on the broad face of lath. The stereographic projection of Fig. 1(b) shows the crystallographic information for this interface. Close packed planes of two phases are misoriented by 1.0° whereas angular deviation of close packed directions is 3.7° for this interface. The Burgers vector of dislocations was determined to be b1 ¼ a=2½0 1 1c ¼ a=2½1 1 1a0 by contrast analysis. Trace analysis revealed that the microscopic line direction of dislocations is ½0:60; 0:57; 0:57a0 on the atomic habit plane ð1 1 1Þckð0 1 1Þa0 , indicating
Fig. 1. (a) Dark-field TEM micrographs of the dislocations (see white lines) on the broad face of lath martensite (223 K, 28.8 ks transformed), (b) ½0 0 1a0 stereographic projection showing the OR, the habit plane and the nature of dislocations for the interface in (a).
T. Moritani et al. / Scripta Materialia 47 (2002) 193–199
that those dislocations are in a pure-screw orientation within an error of 2°. This result is consistent with the observation made by Sandvik and Wayman [6]. In HREM observation, two different beam directions were chosen; the parallel close-packed directions ½1 0 1ck½1 1 1a0 of the K–S OR and ½1 1 0ck½1 0 0a0 which are parallel in the N OR. Fig. 2 shows the HREM images of the broad face of martensite lath observed along these two directions. The broad face of lath, edge-on along ½101ck½111a0 (Fig. 2(a)), contains regularly spaced monoatomic steps with the ð111Þckð011Þa0 terrace as previously reported [8]. These steps can be regarded as transformation dislocations. Along ½110ck½100a0 (Fig. 2(b)), it is seen that many accommodation dislocations with Burgers vectors lying on the parallel close packed planes are present on the interface. At the edge of lath, such a dislocation existed per every third layer of the parallel close packed planes. Fig. 3(a) schematically shows the shear strain originated from the stacking sequence change of
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parallel close packed planes. A strain of a= 12½1 1 2cð¼ a=6½0 1 1a0 Þ, which is perpendicular to ½1 1 0ck½1 0 0a0 , is associated with the stacking change per one ð1 1 1Þc layer. By the coalescence of six transformation dislocations, the shear strain of a=2½1 1 2c is accumulated. This strain can be accommodated by two kinds of perfect dislocations; b1 ¼ a=2½0 1 1cð¼ a=2½1 1 1a0 Þ, b2 ¼ a=2½1 0 1cð¼ a=2½1 1 1aÞ. By introducing these two dislocations alternatively on every third layer of (1 1 1)c, the shear strain due to the stacking change of Fig. 3(a) is fully accommodated. There is another strain associated with the shape change of parallel close packed planes. To accommodate this component, as shown in Fig. 3(b), the same two sets of dislocations b1 and b2 (the solid lines in the figure) need to be introduced on the ð1 1 1Þckð0 1 1Þa0 plane. Although both of those dislocations are almost of pure-screw type on the interface when N OR is held across the boundary, they can be visualized as extra half planes when they are viewed along ½1 1 0ck½1 0 0a0 . Suppose that those dislocations exist as loops on the slip planes inclined on the parallel close packed planes. They can accommodate both of the strain due to the shape change on the parallel close packed planes (i.e., the atomistic habit plane of the broad face) and the strain arising from stacking sequence change at the edge, simultaneously. The dislocations observed in Fig. 2(b) can be explained well by the combination of those two kinds of dislocations. Thus, it is concluded that the shear strain due to the change in the stacking and the shape of parallel close packed planes are fully accommodated by such an arrangement of interphase boundary dislocations. 3.2. Bainitic ferrite/austenite interphase boundary
Fig. 2. HREM micrographs showing the broad face of lath martensite viewed along (a) ½1 0 1ck½1 1 1a0 and (b) ½1 1 0ck½1 0 0a0 , respectively.
Bainitic ferrite formed at 723 K is lath shaped with a thickness of about 1 lm. Each lath consists of smaller sub-units whose thickness is in the order of 0.1 lm, as previously reported [19,20]. Sub-units of ferrite exhibit the ORs scattered around K–S, N and G–T relationships and its habit plane deviates around ð1 2 1Þcðkð1 3 2Þa0 Þ. In the dark-field TEM micrograph of Fig. 4(a), straight dislocations, 5 nm spaced, are present on
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Fig. 3. (a) Stacking sequence change for the parallel close packed planes viewed along ½1 1 0ck½1 0 0a0 and (b) atomic matching on the ð1 1 1Þckð0 1 1Þa0 plane when N OR is held.
Fig. 4. (a) Dark-field TEM micrograph showing the dislocations on the broad face of bainitic ferrite lath (723 K 250 s transformed), (b) ½0 0 1a stereographic projection representing the OR, the habit plane and the nature of dislocations for the interface in (a).
the broad face of bainitic ferrite lath. The stereographic projection of Fig. 4(b) shows that the close packed planes of the two phases are misoriented by 1.5° whereas the angular deviation between the close packed directions is less than 1°. The Burgers vector of dislocations determined by contrast analysis is b1 ¼ a=2½0 1 1c ¼ a=2½1 1 1a0 , which is the same as martensite. Trace analyses of the habit plane and the line direction of dislocations revealed that these dislocations are in a purescrew orientation on the atomic habit plane ð1 1 1Þckð0 1 1Þa0 within an error of less than 2°. The observed dislocation structure was compared with the arrangement of dislocations with no residual strain predicted by O-lattice calculation. In the present study, the same approach as
the study by Hall et al. [21] was used for appropriate lattice parameter in the material used. Fig. 5 shows the calculated dislocation structures on the conjugate parallel close packed plane for ORs varying from the K–S to the N relationship. In the case of N OR (Fig. 5(a)), both sets of accommodation dislocations, are nearly in screw orientations. On the other hand, for the K–S relationship (Fig. 5(c)), the dislocation with b1 ¼ a=2½0 1 1cð¼ a=2½1 1 1aÞ has a mixed character whereas the dislocation with b2 ¼ a=2½1 0 1cð¼ a=2½1 1 1aÞ remains nearly a screw type. For the OR which corresponds to the boundary shown in Fig. 4, the dislocation with b1 is of a mixed type (Fig. 5(b)). The angular deviation between the Burgers vector and the line direction of dislocation D1 is as much
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Fig. 5. O-lattice calculation for the ð1 1 1Þckð0 1 1Þa interface with a misorientation ðhÞ between ½1 0 1c and ½1 1 1a: (a) h ¼ 5:26° (N OR); (b) h ¼ 1° (the OR for Fig. 4); (c) h ¼ 0° (K–S OR). The Burger vectors are b1 ¼ a=2½0 1 1cð¼ a=2½1 1 1aÞ and b2 ¼ a=2½1 0 1cð¼ a=2½1 1 1aÞ.
as 42°. This line direction is significantly different from the observation, implying that misfit strain is not fully accommodated on the ð1 1 1Þckð0 1 1Þa interface. Fig. 6 shows the atomic structures on the broad face of bainitic ferrite. In Fig. 6(a), the macroscopic habit plane consists of monoatomic steps identical to the transformation dislocations observed on the martensite/austenite boundary shown in Fig. 2(a). In Fig. 6(b), accommodation dislocations similar to those on the boundary of lath martensite in Fig. 2(b) are also visible. It is clear that the bainitic ferrite/austenite boundary structure is similar to those of lath martensite with residual transformation strain.
On the contrary to the observation of Fig. 4, sessile bainitic ferrite/austenite boundary structures were reported previously [12,13] in good agreement with the equilibrium dislocation structure predicted by a geometrical model [22]. Some boundaries containing the dislocations in sessile orientations were also recognized in the present study, of which detail is reported elsewhere [23]. Furthermore, it was also proposed [12] that the broad face of bainitic ferrite contains triatomic structural ledges to accommodate misfit in an elastic manner. In the present observation, however, triatomic or higher steps always accompanied some accommodation dislocations.
4. Discussion
Fig. 6. HREM micrographs showing the broad face of bainitic ferrite viewed along (a) ½1 0 1ck½1 1 1a and (b) ½1 1 0ck½1 0 0a.
In the present study, similar pure-screw accommodation dislocations on ð1 1 1Þckð0 1 1Þa were observed in both cases of lath martensite and bainitic ferrite. One of the present authors [24] proposed that bainitic ferrite has a glissile array of interphase boundary dislocations in the early stage. The pure-screw dislocations presently observed on the broad face of bainitic ferrite appears consistent with that proposal but contradicts to the previous study reporting the presence of sessile dislocations on the broad face of bainitic ferrite. The question now arises why both sessile and glissile arrangements of accommodation dislocations are observed on the bainitic ferrite/austenite boundary. The present authors consider that the nature of interphase boundary depends on degrees of strain
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relaxation. There are two strain components in the atomic structure change across the ð1 1 1Þckð0 1 1Þa boundary. One originates from the shape change as described in Fig. 3(b) and the other is the volumetric strain arising from the difference in atomic density on the plane. When interphase boundary diffusion is involved in the stress relaxation around particles, only the strain due to shape change can be relaxed, resulting in a hydrostatic stress field acting on the boundary [25]. When there is a difference in atomic density between the two phases, long-range diffusion is necessary to relax that component. Thus, the transport of extra atoms by volume diffusion of substitutional atoms is required for boundary migration. It should take extremely longer time than the isothermal holding time employed in the present study for bainitic ferrite to grow to the observed sizes. One may consider that the metastable glissile boundary structure changes to the equilibrium sessile structure after sub-units of bainitic ferrite stop growing by some reasons, e.g., accumulation of elastic strain. Such modification process was previously proposed [26–29] although it was regarded that they seem infeasible [30]. Transition of boundary structure should occur by diffusional stress relaxation. If accommodation was aided by short-circuit diffusion, the relaxation time required could be in the order of isothermal transformation time. Since many intruder dislocations are attached to the boundary due to plastic deformation in bainitic ferrite and austenite, fast diffusion along such dislocations might partly affect to the kinetics of strain accommodation. On the transition of interphase boundary structure, however, there can be another point of view which is completely opposite to the one discussed above. Suppose that a sessile boundary is formed during bainitic transformation. During quenching from transformation temperature, the driving force for transformation increases and finally exceeds the critical one necessary for martensitic growth. Under such a circumstance, the equilibrium sessile boundary structure might change to a metastable glissile one. Thus, further consideration is necessary to clarify the detail of strain accommodation process.
Another and more important question to be addressed is whether the boundary observed with two sets of screw dislocations are truly glissile. Christian and Knowles [31] made the proposal that the glissile boundary containing two sets of dislocations are either with two different Burgers vectors on the same glide plane or with the same Burgers vector but on different glide planes. The combination of two sets of screw dislocations in Fig. 3(b) is sessile when their glide planes are different each other as described in the preceding section. If so, the interphase boundary of lath martensite observed becomes sessile. Sandvick and Wayman [7] proposed the operation of the same glide plane for the two sets of dislocations with different Burgers vectors to avoid conflicts. However, the glide planes of those dislocations still need to be identified in future studies. 5. Summary The present study clearly indicates that similar boundary structures can be formed during martensitic and bainite transformations in steels. Acknowledgements The authors are grateful to Professor Hubert I. Aaronson (Carnegie Mellon University, USA) and Professor Gary R. Purdy (McMaster University, Canada) for stimulating discussions. References [1] Clark HM, Wayman CM. In: Aaronson HI, editor. Phase transformations. Metals Park, Ohio: ASM; 1970. p. 59– 114. [2] Frank FC. Acta Metall 1953;1:15–21. [3] Suzuki H. Sci Rep Tohoku Univ 1954;6:30–50. [4] Olson GB, Cohen M. Metall Trans A 1976;7A:1905–14. [5] Sandvik BPJ, Wayman CM. Metall Trans A 1983;14A: 835–44. [6] Sandvik BPJ, Wayman CM. Metall Trans A 1983;14A: 823–34. [7] Sandvik BPJ, Wayman CM. Metall Trans A 1983;14A: 2455–77. [8] Mahon GJ, Howe JM, Mahajan S. Philos Mag Lett 1989;59:273–9.
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