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Nanocrystallization and martensitic transformation in Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy by surface mechanical attrition treatment
Materials Science and Engineering A 481–482 (2008) 484–488
Nanocrystallization and martensitic transformation in Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy by surface mechanical attrition treatment Chunsheng Wen, Wei Li, Yonghua Rong ∗ School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, PR China Received 3 July 2006; received in revised form 15 March 2007; accepted 15 March 2007
Abstract Nanocrystalline grains can be obtained in the surface layer of an Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy with low stacking-fault energy through surface mechanical attrition treatment, accompanying three kinds of strain-induced martensitic transformations. The microstructure of the surface layer was investigated using optical microscopy, X-ray diffraction and transmission electron microscopy. The results indicate the majority of ␣ martensites can be formed directly from the original matrix (␥, fcc), instead of forming at intersections of strain-induced martensites in ␥ matrix grains. The nanocrystallization of grains has three approaches: both the intersection of strain-induced (hcp) martensites and the formation of strain-induced ␣(bcc) martensites from austenite lead to refinement of austenite grains, and the martensitic transformation from (hcp) to ␣(bcc) makes the grain sizes of the product ␣(bcc) smaller than those of (hcp). The strain-induced ␣(bcc) martensites formed from both austenite matrix and (hcp) martensites undergo evolution from dislocation tangles, low angle grain boundaries to large angle grain boundaries. © 2007 Elsevier B.V. All rights reserved. Keywords: Surface mechanical attrition treatment; Strain-induced martensite; Stacking-fault energies
1. Introduction Surface mechanical attrition treatment (SMAT), accomplished by surface shot peening, is demonstrated to be an effective approach to create localized plastic deformation resulting in grain refinement progressively down to the nanometer scale in the surface layer of metallic materials [1]. It has been successfully applied in many material systems [2–6] to achieve surface nanocrystallization. Compared with other surface nanocrystallization methods, SMAT does not result in a difference in chemical compositions between the nanocrystalline surface layer and the matrix, and produces a variation of the grain size from nano-sized (in the top surface layer) to coarse-grains (in the matrix), avoiding bonding problem of the nanocrystalline surface layer with the matrix associated with other methods. In addition, SMAT is simple and low cost, so it has a potential for application in industry.
∗
Corresponding author. Tel.: +86 21 62932558; fax: +86 21 62933910. E-mail address: [email protected] (Y. Rong).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.03.125
It is well known that Fe–Mn–Si-based shape memory alloys are functional materials with low cost and good shape memory effect (SME) [7,8]. A large number of investigations have been published, focusing on martensitic transformations and the improvement of the SME [9–12]. There are normally three kinds of phases ␥(fcc austenite), ␣(bcc martensite), and (hcp martensite) in Fe–Mn–Si-based alloys, and three types of martensitic transformations from ␥(fcc) to (hcp), from ␥(fcc) to ␣(bcc) and (hcp) to ␣(bcc) have been described in previous studies [13–15]. Based on Fujita and Katayama’s [16] observation, ␣(bcc) formed always in the intersection of (hcp) plates. The larger strain can result in more intersection of (hcp), and is favorable to the formation of ␣(bcc). Kajiwara [17] suggested that dislocations do not favor the nucleation of (hcp), instead a high dislocation density results in work hardening and the nucleation of (hcp) will be impeded. In order to make potential use of the SMAT in conventional engineering materials, it is necessary to understand the microstructural features associated with SMAT. In this work, a nanocrystalline surface layer was obtained in the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy (stacking-fault energy (SFE) less than 10 mJ/m2 ) through SMAT. The microstructure
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evolution of this surface region was characterized by different methods. The nanocrystallization mechanism accompanying the martensitic transformation was then proposed. 2. Experimental procedures The chemical composition of a plate of material used in this investigation was determined using energy dispersive analysis of X-rays (EDX) in a scanning electron microscope (SEM515 EDX). The results in wt.% were obtained as 23.4Mn, 6.5Si, 5.1Cr and bal. Fe. The plate was first heated at 1173 K for 60 min followed by water quenching to obtain a uniform grain size. Samples of 70 mm × 70 mm × 5 mm were prepared by SMAT at a frequency of 50 Hz with spherical stainless steel balls of 8 mm diameter. In this work, all the samples were treated in the chamber under vacuum at ambient temperature for different times. The lattice structure of the surface layer was studied by X-ray diffraction (XRD) on a D8 Discover with GADDS X-ray diffractometer with Cr K␣ radiation and the angular range of the diffraction angle (2θ) was between 40◦ and 140◦ . Microstructures of the SMAT samples were examined using optical microscopy (OM) and transmission electron microscopy (TEM; JEM-100CX). Thin foils for TEM investigation were prepared by cutting surface layers with thicknesses of about 0.3 mm, and then mechanically polishing the untreated side until a thickness of 50 m retained. Finally, the TEM foils of 50 m thickness taken out different depths from the top surface were prepared by polishing the treated side. The final thinning was accomplished by twin-jet electropolishing in a solution of 5% perchloric acid and 95% alcohol at about 253 K.
Fig. 2. Optical photograph of Fe–Mn–Si–Cr alloy held at 1173 K for 60 min following by water quenching.
of (hcp). With the increasing treatment duration, the amount of (hcp) first increases and then decreases, while ␣(bcc) appears and its amount gradually increases till a fully ␣(bcc) material was finally obtained. Fig. 2 shows a typical microstructure of the sample heat treated at 1173 K for 60 min followed by water quenching. Thermally induced (hcp) was observed in some grains, and some annealing twins are clearly visible due to the low SFE of this alloy. To study the effect of the inhomogeneous plastic deformation on the microstructure, TEM investigations were carried out on the surface layer of the SMAT sample, as shown in Fig. 3. It was found that the top surface layer consists of roughly equiaxed nanocrystalline grains, and the selected-area electron diffraction (SAED) pattern inserted in indicates that these nanocrystalline grains consist of bcc martensite with random crystallographic orientations.
3. Results and discussion 3.2. Nanocrystallization mechanism 3.1. Microstructure of the surface layer Fig. 1 shows XRD patterns of the quenched sample and SMAT (10 min, 30 min, 60 min and 90 min) samples for the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy. It can be clearly seen that the quenched sample consists of ␥(fcc) and a small amount
Fig. 1. XRD spectra of the surface layer in Fe–Mn–Si–Cr alloy before and after SMAT.
Based on the microstructural features observed in surface regions at different depths of the 60 min SMAT Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy, the nanocrystallization mechanism was revealed as follows: Firstly, lots of stacking faults and strain-induced martensites intersect each other in the austenite, refining austenite grains (Fig. 4). The density of dislocations increases in the austenite
Fig. 3. Dark field image and corresponding diffraction pattern of the nanocrystalline grains in Fe–Mn–Si–Cr alloy after 60 min SMAT.
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Fig. 4. Bright-field image of intersection of stacking faults in Fe–Mn–Si–Cr alloy after 60 min SMAT.
grains with increasing SMAT exposure. Secondly, the straininduced (hcp) grains are refined during further SMAT exposure, as shown in Fig. 5. From the bright field image in Fig. 5(a), it can be concluded that (hcp) plates are divided into three grains (marked with “A”, “B” and “C”) with a substructure in them. Comparing with the dark field images in Fig. 5(c and d), the intersection place of the (hcp) plates can become the nucleation site of ␣(bcc), and thus the → ␣ transformation can result in a (hcp) grain refinement. Finally, the strain-induced ␣(bcc) grains are refined with increasing SMAT exposure, as shown in Fig. 6. Fig. 6 exhibits the typical morphology of ␣(bcc) in the surface layer of this alloy after SMAT. It is clear that the ␣(bcc) is composed of two subgrains in Fig. 7a (marked as “A” and “B” respectively), the corresponding dark field images in Fig. 6(c and d) is obtained from different diffraction spots in [0 1 1]␣ (Fig. 6b). In larger grains of the ␣(bcc) a great number of dislocation tangling can be seen in each subgrain. It can be inferred that an ␣(bcc) forms by a transformation from ␥(fcc) parent phase, not from an intersection of (hcp). With a further severe
plastic deformation, the highly-misoriented grain boundaries in ␣(bcc) grains form gradually through the evolution of low angle boundaries of subgrains, verified by the polycrystal diffraction rings shown in Fig. 7, finally the ␣(bcc) nanocrystalline grains form (Fig. 3). This nanocrystallization mechanism revealed for the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) is somewhat different from that observed for other metallic alloys [18,19]. For example, the SFE of the alloy is smaller than that of an Fe–30Ni (wt.%) alloy (19–35 mJ/m2 ), so there exist ␥ → , → ␣ and ␥ → ␣ martensitic transformations in the Fe–23.4Mn–6.5Si–5.1Cr alloy, different from only the ␥ → ␣ martensitic transformation in the Fe–30Ni (wt.%) alloy. As a result, ␥ → martensitic transformation also is partly responsible for the refinement of austenite grains in this alloy. 3.3. Martensitic transformations during SMAT During SMAT of the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy, nanocrystalline grains can be obtained in the surface layer accompanying strain-induced martensitic transformation. According to the results reported above, it is worthy to note that the majority of the strain-induced martensite is ␣(bcc) rather than (hcp). (hcp) martensite forms only in the early SMAT stages corresponding to small surface strains. The reverse transformation of stress-induced (hcp) martensite in Fe–Mn–Si-based alloys is the basis for its shape memory effect, while the strain-induced ␣(bcc) martensite or thermal-induced (hcp) martensite are unfavorable to this shape memory effect. The formation of (hcp) occurs by glide of a/6<1 1 2>Shockley partial dislocations in every other {1 1 1} plane in austenite. The mechanism of the ␥(fcc) → (hcp) martensitic transformation in alloys with low stacking-fault energy has long been debated. The pole mechanism [20] is not
Fig. 5. Morphology of (hcp) martensites in 60 min SMAT Fe–Mn–Si–Cr alloy (a) bright field image, (b) diffraction pattern, [2 1¯ 1¯ 0] //[0 1 1]␥ , (c) g = (0 1¯ 1 1) , (d) g = (0 1¯ 1 0) .
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Fig. 6. The subgrains of ␣(bcc) martensite in 60 min SMAT Fe–Mn–Si–Cr alloy (a) bright field image, (b) selected-area diffraction pattern, B = [0 1 1]␣ , (c) g = (2 0 0)␣ , (d) g = (0 1¯ 1)␣ .
universally accepted due to its inconsistence with many experiments [7,21] although it gave a clear picture of the nucleation. Hsu [22,23] proposed that thermal- and stress-induced (hcp) in such alloys nucleate by stacking fault going from irregular stacking to regular stacking. The mechanism is described as the formation of faults through the dissociation of one a/21 1 0 perfect dislocation into Shockley partial dislocations of a/62 1 1 on every other {1 1 1} plane in the fcc structure. In this case of regular overlapping of stacking faults, the parent ␥(fcc) can transform to a perfect (hcp) martensite, if some irregular stack-
ing appears, the corresponding defects must remain within the formed martensite. Before that, Olson and Cohen [24] had emphasized that the stacking faults play an important role in the nucleation of (hcp) martensite. During SMAT in the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy, a great deal of dislocations are generated and accumulate in parent ␥(fcc), which favor the formation of (hcp) martensite from the viewpoint of the pole mechanism. But they are unfavorable for the formation of (hcp) martensite from the viewpoint of regular stacking of stacking faults. The results from this work indicate
Fig. 7. The ␣(bcc) martensite with highly-misoriented grain boundaries in 60 min SMAT Fe–Mn–Si–Cr alloy (a) bright-field image, (b) selected-area diffraction pattern, (c) dark field image, g = (2 0 0)␣ .
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that the stacking faults in the alloy can hardly arrange regularly to form (hcp) in such complicated stress fields except in the early stages of SMAT, i.e., a large density of dislocations counteract the nucleation and growth of (hcp) during the treatment. As a consequence, the majority of nanocrystalline ␣(bcc) forms as a result of the ␥ → ␣ transformation, rather than of ␥ → → ␣ transformations, although the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy has low SFE, the ␥ → martensitic transformation does not occur in the presence of severe plastic deformation. 4. Conclusions (i) Nanocrystalline grains can be obtained in the surface layer of the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy with low stacking-fault energy through a surface mechanical attrition treatment, enforcing two kinds of straininduced martensitic transformations. The majority of ␣(bcc) martensites with dislocation microstructure formed directly from the original matrix (␥, fcc), instead of forming at intersections of strain-induced (hcp) martensites. (ii) The refinement of grains occurs due to three approaches: both the intersection of strain-induced (hcp) martensites and the formation of strain-induced ␣(bcc) martensites from austenite lead to refinement of austenite grains, and the martensitic transformation from (hcp) to ␣(bcc) makes the grain sizes of the product ␣(bcc) smaller than those of (hcp). The strain-induced ␣(bcc) martensite from both austenite matrix and (hcp) martensites undergoes evolution from dislocation tangles, low angle grain boundaries to large angle grain boundaries, from which nanocrystalline ␣(bcc) martensite forms in the surface layer. (iii) Under severe plastic deformation conditions, the majority of nanocrystalline ␣(bcc) is a result of a ␥ → ␣ transformation, rather than that of ␥ → → ␣ transformations in
the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy with low SFE, which supports the stacking-fault mechanism of martensite nucleation. References [1] [2] [3] [4] [5] [6] [7] [8] [9]
[10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23]
[24]
K. Lu, J. L¨u, J. Mater. Sci. Technol. 15 (1999) 193. N.R. Tao, M.L. Sui, J. L¨u, K. Lu, Nanostruct. Mater. 11 (1999) 433. X.P. Yong, G. Liu, J. L¨u, K. Lu, Acta Metall. Sin. 38 (2) (2002) 157. G. Liu, J. L¨u, K. Lu, Mater. Sci. Eng. A 286 (2000) 91. N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. L¨u, K. Lu, Acta Mater. 50 (2002) 4603. X. Wu, N. Tao, Y. Hong, B. Xu, J. L¨u, K. Lu, Acta Mater. 50 (2002) 2075. S. Kajiwara, Mater. Sci. Eng. A 273–275 (1999) 67. T.W. Duerig, Mater. Sci. Forum 156–58 (1990) 679. H. Otsuka, M. Murakami, S. Matsuda, Proceedings of the MRS International Meetings on Advanced Materials Shape Memory Materials, vol. 9, Tokyo, 1988, MRS, 1989, p. 451. K. Yamaguchi, Y. Morioka, Y. Tomota, Scripta Mater. 35 (1996) 1147. K. Tsuzaki, Y. Murakami, Y. Natsume, T. Maki, Trans. Mater. Res. Soc. Jpn. 18B (1994) 961. Q. Gu, J.V. Humbeek, L. Delaey, Scripta Metall. 30 (1994) 1587. A. Sato, K. Soma, T. Mori, Acta Metall. 30 (1982) 1901. S. Takaki, H. Nakatsu, Y. Tokunaga, Mater. Trans. JIM 34 (1993) 489. B.X. Huang, X.D. Wang, Y.H. Rong, L. Wang, L. Jin, Mater. Sci. Eng. A 438–440 (2006) 306. H. Fujita, T. Katayama, Mater. Trans. JIM 33 (3) (1992) 243. S. Kajiwara, Metall. Trans. A 17 (1986) 1693. K. Lu, J. L¨u, Mater. Sci. Eng. A 375–377 (2004) 38. C.S. Wen, Z. Chen, B.X. Huang, Y.H. Rong, Metall. Mater. Trans. A 37 (2006) 1413. A. Seeger, Z. Metallk. 47 (1956) 653. T. Kikuchi, S. Kajiwara, Y. Yomota, Mater. Trans. JIM 36 (1995) 719. T.Y. Hsu, Sci. China (in Chinese), Scr. E 27 (4) (1997) 289. Y.H. Rong, Z.H. Guo, S.P. Chen, T.Y. Hsu, Y.X. Guo, in: Koiwa, K. Otsuka, T. Miyazaki (Eds.), Proceedings of the International Conference on Solid–Solid Transformations ’99 (JIMIC-3), The Japan Institute of Metals, Kyoto, 1999, p. 1024. G.B. Olson, M. Cohen, Metall. Trans. A 7 (1976) 1897.
Nanocrystallization and martensitic transformation in Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy by surface mechanical attrition treatment Chunsheng Wen, Wei Li, Yonghua Rong ∗ School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, PR China Received 3 July 2006; received in revised form 15 March 2007; accepted 15 March 2007
Abstract Nanocrystalline grains can be obtained in the surface layer of an Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy with low stacking-fault energy through surface mechanical attrition treatment, accompanying three kinds of strain-induced martensitic transformations. The microstructure of the surface layer was investigated using optical microscopy, X-ray diffraction and transmission electron microscopy. The results indicate the majority of ␣ martensites can be formed directly from the original matrix (␥, fcc), instead of forming at intersections of strain-induced martensites in ␥ matrix grains. The nanocrystallization of grains has three approaches: both the intersection of strain-induced (hcp) martensites and the formation of strain-induced ␣(bcc) martensites from austenite lead to refinement of austenite grains, and the martensitic transformation from (hcp) to ␣(bcc) makes the grain sizes of the product ␣(bcc) smaller than those of (hcp). The strain-induced ␣(bcc) martensites formed from both austenite matrix and (hcp) martensites undergo evolution from dislocation tangles, low angle grain boundaries to large angle grain boundaries. © 2007 Elsevier B.V. All rights reserved. Keywords: Surface mechanical attrition treatment; Strain-induced martensite; Stacking-fault energies
1. Introduction Surface mechanical attrition treatment (SMAT), accomplished by surface shot peening, is demonstrated to be an effective approach to create localized plastic deformation resulting in grain refinement progressively down to the nanometer scale in the surface layer of metallic materials [1]. It has been successfully applied in many material systems [2–6] to achieve surface nanocrystallization. Compared with other surface nanocrystallization methods, SMAT does not result in a difference in chemical compositions between the nanocrystalline surface layer and the matrix, and produces a variation of the grain size from nano-sized (in the top surface layer) to coarse-grains (in the matrix), avoiding bonding problem of the nanocrystalline surface layer with the matrix associated with other methods. In addition, SMAT is simple and low cost, so it has a potential for application in industry.
∗
Corresponding author. Tel.: +86 21 62932558; fax: +86 21 62933910. E-mail address: [email protected] (Y. Rong).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.03.125
It is well known that Fe–Mn–Si-based shape memory alloys are functional materials with low cost and good shape memory effect (SME) [7,8]. A large number of investigations have been published, focusing on martensitic transformations and the improvement of the SME [9–12]. There are normally three kinds of phases ␥(fcc austenite), ␣(bcc martensite), and (hcp martensite) in Fe–Mn–Si-based alloys, and three types of martensitic transformations from ␥(fcc) to (hcp), from ␥(fcc) to ␣(bcc) and (hcp) to ␣(bcc) have been described in previous studies [13–15]. Based on Fujita and Katayama’s [16] observation, ␣(bcc) formed always in the intersection of (hcp) plates. The larger strain can result in more intersection of (hcp), and is favorable to the formation of ␣(bcc). Kajiwara [17] suggested that dislocations do not favor the nucleation of (hcp), instead a high dislocation density results in work hardening and the nucleation of (hcp) will be impeded. In order to make potential use of the SMAT in conventional engineering materials, it is necessary to understand the microstructural features associated with SMAT. In this work, a nanocrystalline surface layer was obtained in the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy (stacking-fault energy (SFE) less than 10 mJ/m2 ) through SMAT. The microstructure
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evolution of this surface region was characterized by different methods. The nanocrystallization mechanism accompanying the martensitic transformation was then proposed. 2. Experimental procedures The chemical composition of a plate of material used in this investigation was determined using energy dispersive analysis of X-rays (EDX) in a scanning electron microscope (SEM515 EDX). The results in wt.% were obtained as 23.4Mn, 6.5Si, 5.1Cr and bal. Fe. The plate was first heated at 1173 K for 60 min followed by water quenching to obtain a uniform grain size. Samples of 70 mm × 70 mm × 5 mm were prepared by SMAT at a frequency of 50 Hz with spherical stainless steel balls of 8 mm diameter. In this work, all the samples were treated in the chamber under vacuum at ambient temperature for different times. The lattice structure of the surface layer was studied by X-ray diffraction (XRD) on a D8 Discover with GADDS X-ray diffractometer with Cr K␣ radiation and the angular range of the diffraction angle (2θ) was between 40◦ and 140◦ . Microstructures of the SMAT samples were examined using optical microscopy (OM) and transmission electron microscopy (TEM; JEM-100CX). Thin foils for TEM investigation were prepared by cutting surface layers with thicknesses of about 0.3 mm, and then mechanically polishing the untreated side until a thickness of 50 m retained. Finally, the TEM foils of 50 m thickness taken out different depths from the top surface were prepared by polishing the treated side. The final thinning was accomplished by twin-jet electropolishing in a solution of 5% perchloric acid and 95% alcohol at about 253 K.
Fig. 2. Optical photograph of Fe–Mn–Si–Cr alloy held at 1173 K for 60 min following by water quenching.
of (hcp). With the increasing treatment duration, the amount of (hcp) first increases and then decreases, while ␣(bcc) appears and its amount gradually increases till a fully ␣(bcc) material was finally obtained. Fig. 2 shows a typical microstructure of the sample heat treated at 1173 K for 60 min followed by water quenching. Thermally induced (hcp) was observed in some grains, and some annealing twins are clearly visible due to the low SFE of this alloy. To study the effect of the inhomogeneous plastic deformation on the microstructure, TEM investigations were carried out on the surface layer of the SMAT sample, as shown in Fig. 3. It was found that the top surface layer consists of roughly equiaxed nanocrystalline grains, and the selected-area electron diffraction (SAED) pattern inserted in indicates that these nanocrystalline grains consist of bcc martensite with random crystallographic orientations.
3. Results and discussion 3.2. Nanocrystallization mechanism 3.1. Microstructure of the surface layer Fig. 1 shows XRD patterns of the quenched sample and SMAT (10 min, 30 min, 60 min and 90 min) samples for the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy. It can be clearly seen that the quenched sample consists of ␥(fcc) and a small amount
Fig. 1. XRD spectra of the surface layer in Fe–Mn–Si–Cr alloy before and after SMAT.
Based on the microstructural features observed in surface regions at different depths of the 60 min SMAT Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy, the nanocrystallization mechanism was revealed as follows: Firstly, lots of stacking faults and strain-induced martensites intersect each other in the austenite, refining austenite grains (Fig. 4). The density of dislocations increases in the austenite
Fig. 3. Dark field image and corresponding diffraction pattern of the nanocrystalline grains in Fe–Mn–Si–Cr alloy after 60 min SMAT.
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Fig. 4. Bright-field image of intersection of stacking faults in Fe–Mn–Si–Cr alloy after 60 min SMAT.
grains with increasing SMAT exposure. Secondly, the straininduced (hcp) grains are refined during further SMAT exposure, as shown in Fig. 5. From the bright field image in Fig. 5(a), it can be concluded that (hcp) plates are divided into three grains (marked with “A”, “B” and “C”) with a substructure in them. Comparing with the dark field images in Fig. 5(c and d), the intersection place of the (hcp) plates can become the nucleation site of ␣(bcc), and thus the → ␣ transformation can result in a (hcp) grain refinement. Finally, the strain-induced ␣(bcc) grains are refined with increasing SMAT exposure, as shown in Fig. 6. Fig. 6 exhibits the typical morphology of ␣(bcc) in the surface layer of this alloy after SMAT. It is clear that the ␣(bcc) is composed of two subgrains in Fig. 7a (marked as “A” and “B” respectively), the corresponding dark field images in Fig. 6(c and d) is obtained from different diffraction spots in [0 1 1]␣ (Fig. 6b). In larger grains of the ␣(bcc) a great number of dislocation tangling can be seen in each subgrain. It can be inferred that an ␣(bcc) forms by a transformation from ␥(fcc) parent phase, not from an intersection of (hcp). With a further severe
plastic deformation, the highly-misoriented grain boundaries in ␣(bcc) grains form gradually through the evolution of low angle boundaries of subgrains, verified by the polycrystal diffraction rings shown in Fig. 7, finally the ␣(bcc) nanocrystalline grains form (Fig. 3). This nanocrystallization mechanism revealed for the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) is somewhat different from that observed for other metallic alloys [18,19]. For example, the SFE of the alloy is smaller than that of an Fe–30Ni (wt.%) alloy (19–35 mJ/m2 ), so there exist ␥ → , → ␣ and ␥ → ␣ martensitic transformations in the Fe–23.4Mn–6.5Si–5.1Cr alloy, different from only the ␥ → ␣ martensitic transformation in the Fe–30Ni (wt.%) alloy. As a result, ␥ → martensitic transformation also is partly responsible for the refinement of austenite grains in this alloy. 3.3. Martensitic transformations during SMAT During SMAT of the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy, nanocrystalline grains can be obtained in the surface layer accompanying strain-induced martensitic transformation. According to the results reported above, it is worthy to note that the majority of the strain-induced martensite is ␣(bcc) rather than (hcp). (hcp) martensite forms only in the early SMAT stages corresponding to small surface strains. The reverse transformation of stress-induced (hcp) martensite in Fe–Mn–Si-based alloys is the basis for its shape memory effect, while the strain-induced ␣(bcc) martensite or thermal-induced (hcp) martensite are unfavorable to this shape memory effect. The formation of (hcp) occurs by glide of a/6<1 1 2>Shockley partial dislocations in every other {1 1 1} plane in austenite. The mechanism of the ␥(fcc) → (hcp) martensitic transformation in alloys with low stacking-fault energy has long been debated. The pole mechanism [20] is not
Fig. 5. Morphology of (hcp) martensites in 60 min SMAT Fe–Mn–Si–Cr alloy (a) bright field image, (b) diffraction pattern, [2 1¯ 1¯ 0] //[0 1 1]␥ , (c) g = (0 1¯ 1 1) , (d) g = (0 1¯ 1 0) .
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Fig. 6. The subgrains of ␣(bcc) martensite in 60 min SMAT Fe–Mn–Si–Cr alloy (a) bright field image, (b) selected-area diffraction pattern, B = [0 1 1]␣ , (c) g = (2 0 0)␣ , (d) g = (0 1¯ 1)␣ .
universally accepted due to its inconsistence with many experiments [7,21] although it gave a clear picture of the nucleation. Hsu [22,23] proposed that thermal- and stress-induced (hcp) in such alloys nucleate by stacking fault going from irregular stacking to regular stacking. The mechanism is described as the formation of faults through the dissociation of one a/21 1 0 perfect dislocation into Shockley partial dislocations of a/62 1 1 on every other {1 1 1} plane in the fcc structure. In this case of regular overlapping of stacking faults, the parent ␥(fcc) can transform to a perfect (hcp) martensite, if some irregular stack-
ing appears, the corresponding defects must remain within the formed martensite. Before that, Olson and Cohen [24] had emphasized that the stacking faults play an important role in the nucleation of (hcp) martensite. During SMAT in the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy, a great deal of dislocations are generated and accumulate in parent ␥(fcc), which favor the formation of (hcp) martensite from the viewpoint of the pole mechanism. But they are unfavorable for the formation of (hcp) martensite from the viewpoint of regular stacking of stacking faults. The results from this work indicate
Fig. 7. The ␣(bcc) martensite with highly-misoriented grain boundaries in 60 min SMAT Fe–Mn–Si–Cr alloy (a) bright-field image, (b) selected-area diffraction pattern, (c) dark field image, g = (2 0 0)␣ .
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that the stacking faults in the alloy can hardly arrange regularly to form (hcp) in such complicated stress fields except in the early stages of SMAT, i.e., a large density of dislocations counteract the nucleation and growth of (hcp) during the treatment. As a consequence, the majority of nanocrystalline ␣(bcc) forms as a result of the ␥ → ␣ transformation, rather than of ␥ → → ␣ transformations, although the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy has low SFE, the ␥ → martensitic transformation does not occur in the presence of severe plastic deformation. 4. Conclusions (i) Nanocrystalline grains can be obtained in the surface layer of the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy with low stacking-fault energy through a surface mechanical attrition treatment, enforcing two kinds of straininduced martensitic transformations. The majority of ␣(bcc) martensites with dislocation microstructure formed directly from the original matrix (␥, fcc), instead of forming at intersections of strain-induced (hcp) martensites. (ii) The refinement of grains occurs due to three approaches: both the intersection of strain-induced (hcp) martensites and the formation of strain-induced ␣(bcc) martensites from austenite lead to refinement of austenite grains, and the martensitic transformation from (hcp) to ␣(bcc) makes the grain sizes of the product ␣(bcc) smaller than those of (hcp). The strain-induced ␣(bcc) martensite from both austenite matrix and (hcp) martensites undergoes evolution from dislocation tangles, low angle grain boundaries to large angle grain boundaries, from which nanocrystalline ␣(bcc) martensite forms in the surface layer. (iii) Under severe plastic deformation conditions, the majority of nanocrystalline ␣(bcc) is a result of a ␥ → ␣ transformation, rather than that of ␥ → → ␣ transformations in
the Fe–23.4Mn–6.5Si–5.1Cr (wt.%) alloy with low SFE, which supports the stacking-fault mechanism of martensite nucleation. References [1] [2] [3] [4] [5] [6] [7] [8] [9]
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