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Reverse transformation mechanism of martensite to austenite in a metastable austenitic alloy
Materials Science and Engineering A 515 (2009) 32–37
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Reverse transformation mechanism of martensite to austenite in a metastable austenitic alloy Seok-Jae Lee a , Yong-Min Park b , Young-Kook Lee b,∗ a b
Research Institute of Iron and Steel Technology, Yonsei University, Seoul 120-749, Republic of Korea Department of Materials Science and Engineering, Yonsei University, Seoul 120-749, Republic of Korea
a r t i c l e
i n f o
Article history: Received 17 January 2009 Received in revised form 9 February 2009 Accepted 10 February 2009 Keywords: Reverse transformation Strain-induced martensite Diffusive transformation Diffusionless transformation Metastable austenitic alloy
a b s t r a c t The reverse transformation of martensite to austenite in a metastable austenitic alloy was investigated during continuous heating followed by isothermal holding. The diffusionless reverse transformation occurred irrespective of heating rate during continuous heating, resulting in lath-shaped austenite with high dislocation density. During isothermal holding, the equiaxed grains are nucleated and grow due to the diffusive reverse transformation. With further holding time, the diffusionlessly formed austenite laths were recovered to be subgrains. These results are summarized by the reverse transformation–temperature–time diagram. © 2009 Elsevier B.V. All rights reserved.
1. Introduction As one of efforts for the simultaneous improvement of strength and ductility in ultrafine-grained (UFG) alloys [1–18], the grain refining process using the strain-induced martensite and its reverse transformations (called SIMRT process [13]) in metastable austenitic alloys [10–18] has been attractive recently. The SIMRT process can easily make UFG alloys by a single cycle of conventional cold rolling and annealing without severe plastic deformation [5,6], which have an excellent combination of both uniform elongation and strength. The important factors for fabrication of the UFG alloys using the SIMRT process are alloy composition, thickness reduction of cold rolling, critical temperature for complete reverse transformation, repetition of the process, etc. However, the core of the SIMRT process is undoubtedly the complete reverse transformation of strain-induced martensite to austenite. Accordingly, the clear understanding of the reverse transformation is inevitable to the SIMRT process. Takaki and co-workers [10] reported that the reverse transformation of martensite to austenite occurs by diffusionless mechanism at the low ratio of Cr to Ni, whereas diffusive reverse transformation happens at the high ratio of Cr to Ni in metastable Fe–Cr–Ni ternary alloys. Lee and Kwon [11,12] reported that the addition of carbon to the metastable austenitic Fe–Cr–Ni alloys
∗ Corresponding author. Tel.: +82 2 2123 2831; fax: +82 2 312 5375. E-mail address: [email protected] (Y.-K. Lee). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.02.010
changes the reverse transformation mechanism from diffusionless to diffusive transformation because the solute carbon in martensite matrix requires the larger driving force for the diffusionless shear reverse transformation [19]. They also investigated the effect of the heating rate on the reverse transformation behavior and found that the reverse transformation mechanism changes from diffusive to diffusionless with increasing the heating rate in the same metastable austenitic Fe–Cr–Ni alloy. While the equiaxed ultrafine grains with a low dislocation density were observed in the slowly heated specimen, which is a characteristic of diffusive transformation, the lath-shaped reversed austenite with a high dislocation density was observed in the fast heated specimen, indicative of diffusionless transformation [11,12]. Takaki and co-workers [10] proposed the reverse transformation–temperature–time (RTT) diagrams based on the isothermal annealing of two kinds of Fe–Cr–Ni alloys. The RTT diagram shows the temperature range of diffusionless transformation is moved depending on chemical composition implying that the reverse transformation mechanism can be different even at the identical annealing temperature depending on alloy chemistry. However, there are little comments about the reverse transformation mechanism of martensite to austenite during the continuous heating followed by isothermal holding. In this study, therefore, the reverse transformation mechanism during the continuous heating and isothermal holding in a metastable austenitic alloy was investigated in aspects of transformation kinetics and microstructural morphology.
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2. Experimental procedure An austenitic alloy, whose chemical composition is Fe–0.023% C–10.91% Cr–9.23% Ni–6.97% Mn in mass percent, was made as an ingot using a high-frequency vacuum induction furnace. The ingot was homogenized at 1200 ◦ C for 12 h and hot-rolled at around 1000 ◦ C to a plate of 10 mm thick. The hot-rolled plate was coldrolled with thickness reduction of 75% to change austenite to fully martensite. The dilatometric specimens of 3 mm × 1 mm × 10 mm were taken from the cold-rolled sheet and were heated up to 1000 ◦ C at various rates of 0.5–100 ◦ C s−1 to measure austenite start (As ) and finish (Af ) temperatures. Some dilatometric specimens were heated up to As + 10 ◦ C at the rate of 10 ◦ C s−1 and held for the maximum 30 min, and then quenched to room temperature by blowing a nitrogen gas in a dilatometer chamber. The volume fractions of austenite and martensite were analyzed using X-ray diffraction (XRD) with Cu K␣ radiation ( = 1.542 Å). The scan speed of the XRD was 2◦ min−1 , the scanned range (2) was between 40◦ and 100◦ , and the voltage and current were 50 kV and 200 mA, respectively. The thin foil specimens for transmission electron microscopy (TEM) were prepared from the dilatometric specimens quenched after different isothermal holding times to observe microstructural changes. The thin foils were jet-polished in a solution of 90% acetic acid + 10% perchloric acid at 15 ◦ C with 15–20 V and 38–40 mA. The jet-polished thin foils were observed using a JEM 2000EX operating at 200 kV. 3. Results and discussion Fig. 1 shows the As and Af temperatures measured at different heating rates using a dilatometer. The As and Af temperatures do not change much regardless of the heating rate, which is known as a characteristic of diffusionless shear transformation [10–12,19–21]. Especially, the Af temperature is very important to determine the lowest temperature for the complete reversed austenite in the SIMRT process. The temperature range of the reverse transformation between As (about 510 ◦ C) and Af temperatures (about 560 ◦ C) is 50 ◦ C for the austenitic alloy used in this study. Takaki and coworkers [10] reported that the diffusionless reverse transformation
Fig. 1. Variations in start (As ) and finish (Af ) temperatures of reverse transformation of martensite to austenite with continuous heating rate.
Fig. 2. Changes in austenite volume fraction and transformation strain during isothermal holding at 520 ◦ C (As + 10 ◦ C). The heating rate up to 520 ◦ C was 10 ◦ C s−1 .
occurred between 530 ◦ C and 630 ◦ C and the transformation temperature range was about 100 ◦ C irrespective of heating rate in Fe–16Cr–10Ni alloy. Lee and Kwon [11,12] also confirmed that the diffusionless reverse transformation occurred between 580 ◦ C and 660 ◦ C and transformation temperature range was about 80 ◦ C in Fe–0.006C–16Cr–9.5Ni alloy. The reason for the decreases in transformation temperature range and As and Af temperatures in the present alloy compared with the previous results [10–12] is higher austenite stability due to more C and Mn contents, especially C content. Fig. 2 shows the variations of austenite volume fraction and transformation strain with isothermal holding time at 520 ◦ C (As + 10 ◦ C). The dilatometric specimen was gradually contracted during isothermal holding, so that the compressive strain (open rectangle) due to the reverse transformation increases linearly with the logarithmic isothermal holding time. The volume fraction of austenite (closed circle) was already about 70% just after the specimen was continuously heated up to 520 ◦ C at 10 ◦ C s−1 . The explosive increase in austenite volume until 520 ◦ C is probably due to the diffusionless reverse transformation from the strain-induced martensite. The volume fraction of austenite linearly increases in proportion to the logarithmic isothermal holding time at 520 ◦ C. With further holding time, the complete reverse transformation is achievable by diffusional transformation. This isothermal transformation is in general a characteristic of time-dependent diffusive transformations. However, some alloys (Fe–Ni, Fe–Ni–Mn) reveal isothermal diffusionless transformation as well as athermal diffusionless transformation during quenching [22]. Machlin and Cohen [23] and Anandaswaroop and Raghavan [24] reported both athermal and isothermal diffusionless martensitic transformations in Fe–29Ni and Fe–30Ni–0.02C alloys, respectively. Therefore, in order to make the reverse transformation mechanism in the present alloy clear, the microstructural evolution during isothermal holding was observed using a TEM. Fig. 3 shows the TEM microstructures of as cold-rolled and isothermally annealed specimens with different holding times at 520 ◦ C. The cold-rolled sample exhibits typical lath martensite whose average lath thickness is about 200 nm (Fig. 3(a)). The cold-rolled specimens were heated up to 520 ◦ C and held for 30 s (Fig. 3(b) and (c)), for 10 min (Fig. 3(d)), and for 30 min (Fig. 3(e)), respectively, and then quenched to room temperature. As shown in Fig. 3(b), the annealed specimen at 520 ◦ C for 30 s shows the lath-shaped microstructure like martensite in Fig. 3(a). However, the selected area diffraction pattern (SADP) taken at the lath boundaries (Fig. 3(c)) proves that the laths are austenite
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Fig. 3. TEM microstructures of (a) martensite in the cold-rolled sample, (b) the annealed specimen at 520 ◦ C for 30 s, and (c) selected area diffraction pattern (SADP) at boundaries of austenite laths in the annealed specimen at 520 ◦ C for 30 s, (d) the annealed specimen at 520 ◦ C for 10 min, and (e) the annealed specimen at 520 ◦ C for 30 min, respectively.
not martensite. This lath-shaped austenite is a characteristic of diffusionlessly transformed austenite, which is also observed in Fe–Cr–Ni alloy [10–12,21], Fe–Ni–C alloy [19], and Fe–Ni alloy [20]. Therefore, at the early stage of the reverse transformation, austenite is formed by diffusionless mechanism. However, dislocation density in the diffusionlessly formed austenite (Fig. 3(b)) seems to be reduced compared with that of the initial cold-rolled martensite. With further annealing time, the new equiaxed austenite grains appear at the untransformed martensite (Fig. 3(d)). They have much
lower dislocation density compared with those of initial martensite (Fig. 3(a)) or diffusionlessly formed austenite (Fig. 3(b)). The austenite grains in Fig. 3(d) show a characteristic of diffusive transformation proceeding by nucleation and growth. The diffusively formed austenite grains gradually grow with increasing holding time by consuming the untransformed martensite, increasing the austenite volume fraction (Fig. 3(e)). Therefore, in the present alloy, the reverse transformation of martensite to austenite first happens diffusionlessly during continuous heating and then additionally
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Fig. 4. Recovery of the diffusionlessly formed austenite in the annealed sample at 520 ◦ C for 30 min. The A, B, and C regions are the subgrains in an austenite lath. (a) TEM microstructure, (b) schematic illustration of the subgrains, and (c) SADPs of three different austenite subgrains.
proceeds in diffusive nucleation and growth manner during isothermal holding. In the meantime, with increasing annealing time, the diffusionlessly formed austenite can be subdivided into subgrains or dislocation cells by recovery process. The subgrains probably change to recrystallized grains by lattice rotation or cell boundary merging [10,25,26]. Fig. 4(a) shows the subgrains in a diffusionlessly formed austenite lath in the annealed specimen at 520 ◦ C for 30 min. As shown in Fig. 4(b), the subgrain boundaries between A and B and between B and C in the austenite lath are still low-angle boundaries with misorientation angles of 3–4◦ . All SADPs of the three different subgrains (A, B, and C) exhibit the same austenite patterns with [0 1 1] zone axis (Fig. 4(c)). The above results show that the recrystallization of the diffusionlessly formed austenite is slower than the diffusive reverse transformation at 520 ◦ C. To obtain a comprehensive understanding about the reverse transformation behaviors of martensite to austenite occurring during continuous heating and isothermal holding in the cold-rolled metastable austenite alloy, the RTT diagram of the alloy used was drawn in Fig. 5. The As and Af temperatures are the measured austenite start and finish temperatures in diffusionless reverse transformation, whereas As and Af temperatures are the expected austenite start and finish temperatures in diffusive reverse transformation, respectively. The dash-dot line indicates the expected recrystallization start (Rs ) curve. The heat treatment schedule con-
Fig. 5. Reverse transformation–temperature–time (RTT) diagram of the metastable austenitic alloy used in this study. As and Af temperatures are the measured diffusionless austenite start and finish temperatures, As and Af temperatures are the expected diffusive austenite start and finish temperatures, respectively. Rs means the expected recrystallization start curve. T0 temperature is an equilibrium transformation temperature where the Gibbs free energies of austenite and ferrite are the same.
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Fig. 6. Schematic illustrations of the microstructural changes during continuous heating and isothermal holding: (a) as cold-rolled, (b) during continuous heating, and (c) during isothermal holding (␣ : martensite; ␥s : diffusionlessly formed austenite; ␥d : diffusively formed austenite). The dash lines in ␥s are subgrain boundaries.
sists of heating at the rate of 10 ◦ C s−1 to As + 10 ◦ C, holding for 30 min at the temperature, and cooling to room temperature. During the continuous heating at 10 ◦ C s−1 to As + 10 ◦ C, the martensite transformed to austenite without atomic diffusion, giving rise to lath-shaped austenite having high dislocation density. When the specimen maintains at the temperature, the diffusionless reverse transformation stops and diffusive reverse transformation immediately starts. Because the transformation strain change due to the diffusive reverse transformation (Fig. 2) immediately appeared without any incubation time as soon as the isothermal holding begins, the As temperature curve is expected to be passed before the isothermal holding at As + 10 ◦ C. This diffusive reverse transformation results in equiaxed austenite grains with low dislocation density. With further holding time, the new equiaxed austenite grains are nucleated and grow and the diffusionlessly formed austenite laths are recovered and subdivided into several subgrains with low-angle grain boundaries. The isothermal holding of 30 min at As + 10 ◦ C is not enough to trigger the recrystallization of diffusionlessly formed austenite laths. Fig. 6 shows schematic illustrations of the microstructural changes during continuous heating and isothermal holding in the cold-rolled metastable austenite alloy as a summary of the abovementioned results. 4. Conclusions In this study, the reverse transformation mechanism of martensite to austenite during continuous heating and isothermal holding
in the cold-rolled metastable austenite alloy was investigated using a dilatometer and a TEM. During the continuous heating, the diffusionless reverse transformation took place irrespective of heating rate, giving rise to lath-shaped austenite, which becomes recovered and subdivided into subgrains with further holding time at As + 10 ◦ C. However, during the isothermal holding at As + 10 ◦ C, the reverse transformation proceeds in a diffusive manner without incubation time, resulting in equiaxed grains, additional austenite volume fraction, and compressive transformation strain. A reverse transformation–temperature–time diagram was proposed based on the microstructural changes during continuous heating and isothermal holding. Acknowledgements This work was supported by the Defense Acquisition Program Administration and the Agency for Defense Development. References [1] [2] [3] [4] [5] [6] [7] [8]
Y.M. Wang, M.W. Chen, F.H. Zhou, E. Ma, Nature 419 (2002) 912–915. R.Z. Valiev, Nature 419 (2002) 887–889. L. Lu, Y. Shen, X. Chen, L. Qian, K. Lu, Science 304 (2004) 422–426. Y.J. Zhao, X.Z. Liao, S. Cheng, E. Ma, Y.T. Zhu, Adv. Mater. 18 (2006) 2280–2283. Y.M. Wang, E. Ma, Mater. Sci. Eng. A 375 (2004) 46–52. Y.M. Wang, E. Ma, Acta Mater. 52 (2004) 1699–1707. G.J. Fan, H. Choo, P.K. Liaw, E.J. Lavernia, Acta Mater. 54 (2006) 1759–1766. M.R. Shankar, B.C. Rao, S. Chandrasekar, W.D. Compton, A.H. King, Scripta Mater. 58 (2008) 675–678. [9] Y.H. Zhao, Z. Horita, T.G. Langdon, Y.T. Zhu, Mater. Sci. Eng. A 474 (2008) 342–347.
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K. Tomimura, S. Takaki, Y. Tokunaga, ISIJ Int. 31 (1991) 1431–1437. Y.K. Lee, O.J. Kwon, J. Kor. Inst. Mat. Mater. 30 (1992) 1317–1325. Y.K. Lee, O.J. Kwon, J. Kor. Inst. Mat. Mater. 31 (1993) 208–215. Y.Q. Ma, J.E. Jin, Y.K. Lee, Mater. Sci. Forum 475–479 (2005) 43–48. Y.Q. Ma, J.E. Jin, Y.K. Lee, Scripta Mater. 52 (2005) 1311–1315. Y.K. Lee, J.E. Jin, Y.Q. Ma, Scripta Mater. 57 (2007) 707–710. E.S. Perdahcıo˘glu, H.J.M. Geijselaers, J. Huétink, Mater. Sci. Eng. A 481–482 (2008) 727–731. [17] E.S. Perdahcıo˘glu, H.J.M. Geijselaersa, M. Groen, Scripta Mater. 58 (2008) 947–950. [18] R.D.K. Misra, B.R. Kumar, M. Somani, P. Karjalainen, Scripta Mater. 59 (2008) 79–82.
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[19] C.A. Apple, G. Krauss, Acta Metall. 20 (1972) 849–856. [20] S. Jana, C.M. Wayman, Trans. Metall. Soc. AIME 239 (1967) 1187– 1193. [21] D.S. Leem, Y.D. Lee, J.H. Jun, C.S. Choi, Scripta Mater. 45 (2001) 767– 772. [22] Z. Nishiyama, in: M. Fine, M. Meshi, C. Wayman (Eds.), Martensitic Transformation, Academic Press, New York, 1978. [23] E.S. Machlin, M. Cohen, Trans. Metall. Soc. AIME 194 (1952) 489–500. [24] A.V. Anandaswaroop, V. Raghavan, Scripta Metall. 3 (1969) 221–224. [25] T. Tsuchiyama, Y. Miyamoto, S. Takaki, ISIJ Int. 41 (2001) 1047–1052. [26] A. Belyakov, T. Sakai, H. Miura, R. Kaibyshev, K. Tsuzaki, Acta Mater. 50 (2002) 1547–1557.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Reverse transformation mechanism of martensite to austenite in a metastable austenitic alloy Seok-Jae Lee a , Yong-Min Park b , Young-Kook Lee b,∗ a b
Research Institute of Iron and Steel Technology, Yonsei University, Seoul 120-749, Republic of Korea Department of Materials Science and Engineering, Yonsei University, Seoul 120-749, Republic of Korea
a r t i c l e
i n f o
Article history: Received 17 January 2009 Received in revised form 9 February 2009 Accepted 10 February 2009 Keywords: Reverse transformation Strain-induced martensite Diffusive transformation Diffusionless transformation Metastable austenitic alloy
a b s t r a c t The reverse transformation of martensite to austenite in a metastable austenitic alloy was investigated during continuous heating followed by isothermal holding. The diffusionless reverse transformation occurred irrespective of heating rate during continuous heating, resulting in lath-shaped austenite with high dislocation density. During isothermal holding, the equiaxed grains are nucleated and grow due to the diffusive reverse transformation. With further holding time, the diffusionlessly formed austenite laths were recovered to be subgrains. These results are summarized by the reverse transformation–temperature–time diagram. © 2009 Elsevier B.V. All rights reserved.
1. Introduction As one of efforts for the simultaneous improvement of strength and ductility in ultrafine-grained (UFG) alloys [1–18], the grain refining process using the strain-induced martensite and its reverse transformations (called SIMRT process [13]) in metastable austenitic alloys [10–18] has been attractive recently. The SIMRT process can easily make UFG alloys by a single cycle of conventional cold rolling and annealing without severe plastic deformation [5,6], which have an excellent combination of both uniform elongation and strength. The important factors for fabrication of the UFG alloys using the SIMRT process are alloy composition, thickness reduction of cold rolling, critical temperature for complete reverse transformation, repetition of the process, etc. However, the core of the SIMRT process is undoubtedly the complete reverse transformation of strain-induced martensite to austenite. Accordingly, the clear understanding of the reverse transformation is inevitable to the SIMRT process. Takaki and co-workers [10] reported that the reverse transformation of martensite to austenite occurs by diffusionless mechanism at the low ratio of Cr to Ni, whereas diffusive reverse transformation happens at the high ratio of Cr to Ni in metastable Fe–Cr–Ni ternary alloys. Lee and Kwon [11,12] reported that the addition of carbon to the metastable austenitic Fe–Cr–Ni alloys
∗ Corresponding author. Tel.: +82 2 2123 2831; fax: +82 2 312 5375. E-mail address: [email protected] (Y.-K. Lee). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.02.010
changes the reverse transformation mechanism from diffusionless to diffusive transformation because the solute carbon in martensite matrix requires the larger driving force for the diffusionless shear reverse transformation [19]. They also investigated the effect of the heating rate on the reverse transformation behavior and found that the reverse transformation mechanism changes from diffusive to diffusionless with increasing the heating rate in the same metastable austenitic Fe–Cr–Ni alloy. While the equiaxed ultrafine grains with a low dislocation density were observed in the slowly heated specimen, which is a characteristic of diffusive transformation, the lath-shaped reversed austenite with a high dislocation density was observed in the fast heated specimen, indicative of diffusionless transformation [11,12]. Takaki and co-workers [10] proposed the reverse transformation–temperature–time (RTT) diagrams based on the isothermal annealing of two kinds of Fe–Cr–Ni alloys. The RTT diagram shows the temperature range of diffusionless transformation is moved depending on chemical composition implying that the reverse transformation mechanism can be different even at the identical annealing temperature depending on alloy chemistry. However, there are little comments about the reverse transformation mechanism of martensite to austenite during the continuous heating followed by isothermal holding. In this study, therefore, the reverse transformation mechanism during the continuous heating and isothermal holding in a metastable austenitic alloy was investigated in aspects of transformation kinetics and microstructural morphology.
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2. Experimental procedure An austenitic alloy, whose chemical composition is Fe–0.023% C–10.91% Cr–9.23% Ni–6.97% Mn in mass percent, was made as an ingot using a high-frequency vacuum induction furnace. The ingot was homogenized at 1200 ◦ C for 12 h and hot-rolled at around 1000 ◦ C to a plate of 10 mm thick. The hot-rolled plate was coldrolled with thickness reduction of 75% to change austenite to fully martensite. The dilatometric specimens of 3 mm × 1 mm × 10 mm were taken from the cold-rolled sheet and were heated up to 1000 ◦ C at various rates of 0.5–100 ◦ C s−1 to measure austenite start (As ) and finish (Af ) temperatures. Some dilatometric specimens were heated up to As + 10 ◦ C at the rate of 10 ◦ C s−1 and held for the maximum 30 min, and then quenched to room temperature by blowing a nitrogen gas in a dilatometer chamber. The volume fractions of austenite and martensite were analyzed using X-ray diffraction (XRD) with Cu K␣ radiation ( = 1.542 Å). The scan speed of the XRD was 2◦ min−1 , the scanned range (2) was between 40◦ and 100◦ , and the voltage and current were 50 kV and 200 mA, respectively. The thin foil specimens for transmission electron microscopy (TEM) were prepared from the dilatometric specimens quenched after different isothermal holding times to observe microstructural changes. The thin foils were jet-polished in a solution of 90% acetic acid + 10% perchloric acid at 15 ◦ C with 15–20 V and 38–40 mA. The jet-polished thin foils were observed using a JEM 2000EX operating at 200 kV. 3. Results and discussion Fig. 1 shows the As and Af temperatures measured at different heating rates using a dilatometer. The As and Af temperatures do not change much regardless of the heating rate, which is known as a characteristic of diffusionless shear transformation [10–12,19–21]. Especially, the Af temperature is very important to determine the lowest temperature for the complete reversed austenite in the SIMRT process. The temperature range of the reverse transformation between As (about 510 ◦ C) and Af temperatures (about 560 ◦ C) is 50 ◦ C for the austenitic alloy used in this study. Takaki and coworkers [10] reported that the diffusionless reverse transformation
Fig. 1. Variations in start (As ) and finish (Af ) temperatures of reverse transformation of martensite to austenite with continuous heating rate.
Fig. 2. Changes in austenite volume fraction and transformation strain during isothermal holding at 520 ◦ C (As + 10 ◦ C). The heating rate up to 520 ◦ C was 10 ◦ C s−1 .
occurred between 530 ◦ C and 630 ◦ C and the transformation temperature range was about 100 ◦ C irrespective of heating rate in Fe–16Cr–10Ni alloy. Lee and Kwon [11,12] also confirmed that the diffusionless reverse transformation occurred between 580 ◦ C and 660 ◦ C and transformation temperature range was about 80 ◦ C in Fe–0.006C–16Cr–9.5Ni alloy. The reason for the decreases in transformation temperature range and As and Af temperatures in the present alloy compared with the previous results [10–12] is higher austenite stability due to more C and Mn contents, especially C content. Fig. 2 shows the variations of austenite volume fraction and transformation strain with isothermal holding time at 520 ◦ C (As + 10 ◦ C). The dilatometric specimen was gradually contracted during isothermal holding, so that the compressive strain (open rectangle) due to the reverse transformation increases linearly with the logarithmic isothermal holding time. The volume fraction of austenite (closed circle) was already about 70% just after the specimen was continuously heated up to 520 ◦ C at 10 ◦ C s−1 . The explosive increase in austenite volume until 520 ◦ C is probably due to the diffusionless reverse transformation from the strain-induced martensite. The volume fraction of austenite linearly increases in proportion to the logarithmic isothermal holding time at 520 ◦ C. With further holding time, the complete reverse transformation is achievable by diffusional transformation. This isothermal transformation is in general a characteristic of time-dependent diffusive transformations. However, some alloys (Fe–Ni, Fe–Ni–Mn) reveal isothermal diffusionless transformation as well as athermal diffusionless transformation during quenching [22]. Machlin and Cohen [23] and Anandaswaroop and Raghavan [24] reported both athermal and isothermal diffusionless martensitic transformations in Fe–29Ni and Fe–30Ni–0.02C alloys, respectively. Therefore, in order to make the reverse transformation mechanism in the present alloy clear, the microstructural evolution during isothermal holding was observed using a TEM. Fig. 3 shows the TEM microstructures of as cold-rolled and isothermally annealed specimens with different holding times at 520 ◦ C. The cold-rolled sample exhibits typical lath martensite whose average lath thickness is about 200 nm (Fig. 3(a)). The cold-rolled specimens were heated up to 520 ◦ C and held for 30 s (Fig. 3(b) and (c)), for 10 min (Fig. 3(d)), and for 30 min (Fig. 3(e)), respectively, and then quenched to room temperature. As shown in Fig. 3(b), the annealed specimen at 520 ◦ C for 30 s shows the lath-shaped microstructure like martensite in Fig. 3(a). However, the selected area diffraction pattern (SADP) taken at the lath boundaries (Fig. 3(c)) proves that the laths are austenite
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Fig. 3. TEM microstructures of (a) martensite in the cold-rolled sample, (b) the annealed specimen at 520 ◦ C for 30 s, and (c) selected area diffraction pattern (SADP) at boundaries of austenite laths in the annealed specimen at 520 ◦ C for 30 s, (d) the annealed specimen at 520 ◦ C for 10 min, and (e) the annealed specimen at 520 ◦ C for 30 min, respectively.
not martensite. This lath-shaped austenite is a characteristic of diffusionlessly transformed austenite, which is also observed in Fe–Cr–Ni alloy [10–12,21], Fe–Ni–C alloy [19], and Fe–Ni alloy [20]. Therefore, at the early stage of the reverse transformation, austenite is formed by diffusionless mechanism. However, dislocation density in the diffusionlessly formed austenite (Fig. 3(b)) seems to be reduced compared with that of the initial cold-rolled martensite. With further annealing time, the new equiaxed austenite grains appear at the untransformed martensite (Fig. 3(d)). They have much
lower dislocation density compared with those of initial martensite (Fig. 3(a)) or diffusionlessly formed austenite (Fig. 3(b)). The austenite grains in Fig. 3(d) show a characteristic of diffusive transformation proceeding by nucleation and growth. The diffusively formed austenite grains gradually grow with increasing holding time by consuming the untransformed martensite, increasing the austenite volume fraction (Fig. 3(e)). Therefore, in the present alloy, the reverse transformation of martensite to austenite first happens diffusionlessly during continuous heating and then additionally
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Fig. 4. Recovery of the diffusionlessly formed austenite in the annealed sample at 520 ◦ C for 30 min. The A, B, and C regions are the subgrains in an austenite lath. (a) TEM microstructure, (b) schematic illustration of the subgrains, and (c) SADPs of three different austenite subgrains.
proceeds in diffusive nucleation and growth manner during isothermal holding. In the meantime, with increasing annealing time, the diffusionlessly formed austenite can be subdivided into subgrains or dislocation cells by recovery process. The subgrains probably change to recrystallized grains by lattice rotation or cell boundary merging [10,25,26]. Fig. 4(a) shows the subgrains in a diffusionlessly formed austenite lath in the annealed specimen at 520 ◦ C for 30 min. As shown in Fig. 4(b), the subgrain boundaries between A and B and between B and C in the austenite lath are still low-angle boundaries with misorientation angles of 3–4◦ . All SADPs of the three different subgrains (A, B, and C) exhibit the same austenite patterns with [0 1 1] zone axis (Fig. 4(c)). The above results show that the recrystallization of the diffusionlessly formed austenite is slower than the diffusive reverse transformation at 520 ◦ C. To obtain a comprehensive understanding about the reverse transformation behaviors of martensite to austenite occurring during continuous heating and isothermal holding in the cold-rolled metastable austenite alloy, the RTT diagram of the alloy used was drawn in Fig. 5. The As and Af temperatures are the measured austenite start and finish temperatures in diffusionless reverse transformation, whereas As and Af temperatures are the expected austenite start and finish temperatures in diffusive reverse transformation, respectively. The dash-dot line indicates the expected recrystallization start (Rs ) curve. The heat treatment schedule con-
Fig. 5. Reverse transformation–temperature–time (RTT) diagram of the metastable austenitic alloy used in this study. As and Af temperatures are the measured diffusionless austenite start and finish temperatures, As and Af temperatures are the expected diffusive austenite start and finish temperatures, respectively. Rs means the expected recrystallization start curve. T0 temperature is an equilibrium transformation temperature where the Gibbs free energies of austenite and ferrite are the same.
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Fig. 6. Schematic illustrations of the microstructural changes during continuous heating and isothermal holding: (a) as cold-rolled, (b) during continuous heating, and (c) during isothermal holding (␣ : martensite; ␥s : diffusionlessly formed austenite; ␥d : diffusively formed austenite). The dash lines in ␥s are subgrain boundaries.
sists of heating at the rate of 10 ◦ C s−1 to As + 10 ◦ C, holding for 30 min at the temperature, and cooling to room temperature. During the continuous heating at 10 ◦ C s−1 to As + 10 ◦ C, the martensite transformed to austenite without atomic diffusion, giving rise to lath-shaped austenite having high dislocation density. When the specimen maintains at the temperature, the diffusionless reverse transformation stops and diffusive reverse transformation immediately starts. Because the transformation strain change due to the diffusive reverse transformation (Fig. 2) immediately appeared without any incubation time as soon as the isothermal holding begins, the As temperature curve is expected to be passed before the isothermal holding at As + 10 ◦ C. This diffusive reverse transformation results in equiaxed austenite grains with low dislocation density. With further holding time, the new equiaxed austenite grains are nucleated and grow and the diffusionlessly formed austenite laths are recovered and subdivided into several subgrains with low-angle grain boundaries. The isothermal holding of 30 min at As + 10 ◦ C is not enough to trigger the recrystallization of diffusionlessly formed austenite laths. Fig. 6 shows schematic illustrations of the microstructural changes during continuous heating and isothermal holding in the cold-rolled metastable austenite alloy as a summary of the abovementioned results. 4. Conclusions In this study, the reverse transformation mechanism of martensite to austenite during continuous heating and isothermal holding
in the cold-rolled metastable austenite alloy was investigated using a dilatometer and a TEM. During the continuous heating, the diffusionless reverse transformation took place irrespective of heating rate, giving rise to lath-shaped austenite, which becomes recovered and subdivided into subgrains with further holding time at As + 10 ◦ C. However, during the isothermal holding at As + 10 ◦ C, the reverse transformation proceeds in a diffusive manner without incubation time, resulting in equiaxed grains, additional austenite volume fraction, and compressive transformation strain. A reverse transformation–temperature–time diagram was proposed based on the microstructural changes during continuous heating and isothermal holding. Acknowledgements This work was supported by the Defense Acquisition Program Administration and the Agency for Defense Development. References [1] [2] [3] [4] [5] [6] [7] [8]
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