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The precipitation of FCC phase from BCC matrix in an Fe–Mn–Al alloy
Materials Science and Engineering A323 (2002) 462– 466 www.elsevier.com/locate/msea
The precipitation of FCC phase from BCC matrix in an Fe–Mn–Al alloy Wei-Chun Cheng *, Hsin-Yu Lin Department of Mechanical Engineering, National Taiwan Uni6ersity of Science and Technology, Taipei 106, Taiwan Received 10 April 2001; received in revised form 15 May 2001
Abstract The Fe-23.0 wt.% Mn-7.4 wt.% Al-0.03 wt.% C alloy, after water quenching following 1 h at 1050 °C in air, has a small amount of austenite distributed discontinuously along the grain boundaries of the ferrite phase. Aging the as-quenched alloy at higher temperatures caused precipitation of the austenite phase from the ferrite matrix preferentially along grain boundaries or within the matrix. The precipitation of the austenite phase within the ferrite matrix preferred the form of Widmansta¨tten side-plates. Electron diffraction studies verify that the austenite phase precipitates from the ferrite matrix. A Kurdjumov– Sachs orientation relationship holds between the BCC matrix and the FCC precipitate. The proportion of the austenite phase is larger for lower temperatures aging or furnace cooling than in the as-quenched condition. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Quenching; Aging; Precipitation; Widmansta¨tten side-plate; Kurdjumov– Sachs orientation relationship
1. Introduction In recent decades Fe – Mn – Al alloys have gained much attention as substitutes for some of the conventional Fe –Ni –Cr stainless steels. Results concerning corrosion resistance and high temperature oxidation resistance have been reported in several publications [1,2]. For the development of corrosion resistance and high temperature oxidation resistance, the phase equilibria of the Fe – Mn – Al ternary system have practical importance. According to Sato’s work [3], in the range of 20–30 wt.% Mn and 4 – 20 wt.% Al at temperature of 1050 °C, Fe –Mn – Al alloys are composed of ferrite and austenite phases, i.e. some of them are single phase, either ferrite or austenite, and others are duplex phase. In common ferrous alloys, depending on the cooling rates for the austenite to ferrite phase transformations, Widmansta¨tten side-plate, massive, and martensite phases have been observed within the original FCC matrix during the high temperature quenching process [4 –9]. Experimental measurements on Widmansta¨tten * Corresponding author. Tel.: + 886-2-2737-6241; fax: + 886-22737-6460. E-mail address: [email protected] (W.-C. Cheng).
side-plate ferrite phase in common ferrous alloys show that the orientation relationships close to the Kurdjumov –Sachs or Nishiyama –Wasserman types are usually found [10]. In contrast to the common ferrous alloys in the phase transformation during cooling from high temperatures, Fe –Mn –Al alloys have gained much attention for the formation of the FCC phase inside the BCC parent phase. It is well known that among Fe –Mn –Al alloys, Mn and Al are FCC and BCC formers, respectively. Higher Mn content causes the higher proportion of the FCC phase to emerge at low temperatures in contrast to the full BCC phase in common ferrous alloys. If Fe –Mn –Al alloys contain low concentration of Al and high concentration of Mn, the full austenite phase could be observed even at room temperature [3]. In this unusual situation, it may seem that BCC (d-Fe) to FCC (g-Fe) phase transformation at high temperatures has been shifted to lower temperatures. But in common ferrous alloys, this kind of phase transformation is inhibited or screened out by the formation of another BCC (a-Fe) at low temperatures. Fortunately, the addition Mn into the ferrous alloy expands the FCC phase to low temperatures and the BCC (d-Fe) to FCC (g-Fe) phase transformation at higher temperatures can be
0921-5093/02/$ - see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 4 9 8 - 8
W.-C. Cheng, H.-Y. Lin / Materials Science and Engineering A323 (2002) 462–466
shifted to room temperature. For BCC to FCC phase transformations in Fe– Mn – Al alloys, Widmansta¨ tten side-plate [11], massive [12], and 18R type martensitic phases [13–15] have also been observed within the original BCC matrix as the samples are quenched rapidly from 1300 °C. As the FCC phase is stable at low temperatures for some of the Fe– Mn – Al alloys, the purpose of the present work is to study the precipitation of the FCC phase within BCC matrix at low temperatures.
2. Experimental procedures Slabs with composition of Fe-23.0 wt.% Mn-7.4 wt.% Al-0.03 wt.% C were initially prepared by air induction melting. 1008 plain carbon steel, electrolytic manganese, and high– purity aluminum were cast into 10-kg ingots. After being homogenized at 1200 °C for 4 h under a protective argon atmosphere, the ingots were hot forged and then cold rolled to a thickness of 2 mm. These were heated at 1050 °C for 1 h in air, waterquenched to room temperature, then aged at higher temperatures for the study. In addition to the above treatment, samples were also heated at 1050 °C for 1 h in air then furnace cooled at a rate of 50 °C/h to room temperature via program control for comparison. Samples in the above heat treatments were sectioned, mechanically polished and etched in a 10% nital solution for light microscopy (LM) observation. Thin foils for transmission electron microscopy (TEM) were obtained by a twin-jet polisher in a 10% HClO4 and 90% ethanol solution at 30 V and 0.1 A cm − 2, examined in a JEOL 2000 FXII STEM and JEOL JEM 2010 TEM operated at 200 kV. Some of the samples were also examined in a RIGAKU DMAX-B X-ray diffractometer with the maximum power of 12 kW.
Fig. 1. Light micrograph (LM) of the Fe–Mn–Al alloy after quenching into room temperature water from 1 h at 1050 °C (A, austenite; F, ferrite).
463
3. Results and discussion Fig. 1 is the light micrograph (LM) of the Fe– Mn – Al alloy after being quenched into room temperature water following 1 h at 1050 °C. According to LM observation, the alloy consists of two phases: the matrix phase and the discrete phase along the grain boundary of the matrix. In an X-ray study, only ferrite and austenite phases are detected and confirmed by TEM study. The matrix phase is ferrite and the discrete phase along the grain boundaries of the matrix is austenite. The result is quite consistent with Sato’s observation in the two-phase region of Fe –Mn –Al alloys [3]. No precipitates are observed in the as-quenched condition by X-ray method, light or electron microscopy in either phase. The evolution of the precipitates during the aging process at higher temperatures was studied by transmission electron microscopy. In the present study, we conducted aging processes for temperatures ranging from 350 to 700 °C. Fig. 2(a) is a TEM bright-field image of the alloy after being aged at 600 °C for 12 h. It reveals the formation of the Widmansta¨tten sideplate precipitates within the ferrite matrix. The Widmansta¨tten precipitates were identified as the austenite phase, and the selected area diffraction patterns taken in these precipitates are shown in Fig. 3. This result was also confirmed by X-ray study in that only ferrite and austenite phases were detectable. The precipitation of the austenite phase preferred the grain boundaries of the ferrite phase and within the matrix, as shown in Fig. 2(b), in which the sample was aged at 540 °C for 24 h. According to the above observations, the phase transformation of the Fe –Mn –Al alloy during aging is opposite to that of conventional steels in that the low-temperature stable crystal structure is FCC rather than BCC. It is well known that orientation relationships between ferrite and austenite always accompany phase transformations in steel systems, for example martensitic transformation in Fe –Ni –C [4–8], and the formation of Widmannsta¨tten side-plates and intragranular phases in plain carbon steels [9]. Such orientation relationships provide information in which helps to understand the mechanisms of the structural changes. In the present Fe –Mn –Al alloy, during the BCC to FCC phase transformation, we found that the FCC Widmansta¨tten side-plate phase also had the orientation relationships with the parent matrix. Fig. 4(a –b) is selected area diffraction patterns (SADPs) of the FCC phase superimposed on that of the matrix. As indicated in Fig. 4(a –b), the orientation relationship between the FCC phase and the BCC matrix can be derived as: (111)FCC//(110)BCC, [101]FCC//[111]BCC,
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Fig. 2. TEM bright-field images of the alloy after being aged at 600 °C for 12 h (a) and at 540 °C for 24 h (b), respectively. (F, ferrite; G.B., grain boundary).
Fig. 3. The selected area diffraction patterns (SADPs) taken in the precipitates shown in Fig. 2 (a). Zone axes are: (a) [100]; (b) [111]; and (c) [011], respectively (T, transmitted beam).
[110]FCC near [100]BCC, which are the well known Kurdjumov– Sachs (K–S) orientation relationships [13]. This result is quite similar to that in the BCC to FCC phase transformations in Fe –Mn –Al alloys, Widmansta¨ tten side-plate [11], massive [12], and 18R type martensitic phases [13] which also have the same K– S orientation relationships with the original BCC matrix as the samples are quenched rapidly from 1300 °C. Aging the alloy at temperatures below 430 °C, caused the precipitation of D03 phase by homogeneous nucleation and growth and was first observed in the ferrite grains. For longer aging times, some Widmansta¨ tten precipitates were also present in addition to the D03 phase within the ferrite matrix. But at this stage they were too small to be identified. For aging temperatures above 430 °C, no D03 precipitates could
be observed, but the Widmansta¨ tten precipitates were observed under all circumstances. The precipitation of the FCC phase from the BCC matrix was observed during aging after water-quenching the slabs from 1 h at 1050 °C. It is clear that the proportion of the FCC phase was larger at lower temperatures than that at 1050 °C. In order to study the near equilibrium condition for the co-existence of the FCC and BCC phases at lower temperatures, we performed an experiment in which the slabs were heated at 1050 °C for 1 h and then cooled at a rate of 50 °C h − 1 to room temperature. Fig. 5 is the light micrograph (LM) of the Fe–Mn –Al alloy after being heated at 1050 °C for 1 h and then cooled at a rate of 50 °C h − 1 to room temperature. According to the LM observation, the FCC and BCC phases in the alloy are quite different from the condition shown in Fig. 1, and the proportion of the FCC phase is larger than that in
W.-C. Cheng, H.-Y. Lin / Materials Science and Engineering A323 (2002) 462–466
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Fig. 4. SADPs of the FCC phase superimposed on that of the BCC matrix in Fig. 2(b) showing the K – S orientation relationships. (a) Zone axes are [101]FCC and [111]BCC. (b) [110]FCC zone axis is near [100]BCC zone axis. (hkl, FCC; hkl, BCC).
the as-quenched condition. This observation demonstrates that the FCC phase is more stable than the BCC phase at low temperatures. In TEM study, D03 phase has formed within the ferrite phase during program-controlled furnace cooling as shown in Fig. 6(a), which is the g = 200 dark-field image of the D03 phase. Fig. 6(b) is the selected area diffraction pattern taken from the mixed region of ferrite matrix and the D03 fine precipitates. The zone axis is [011]. In addition to the spots corresponding to the ferrite matrix, the diffraction pattern also consists of small superlattice spots of 111, 200 etc. caused by the presence of the precipitates. The formation of the D03 phase within the ferrite phase is quite consistent with the result derived from the aging processes below 430 °C.
[110]FCC near [100]BCC, which correspond to the K–S orientation relationships. It is clear that the proportion of the austenite phase in the two-phase region of the Fe–Mn –Al alloy is larger at lower temperatures during aging or furnace cooling than that for the as-quenched condition from 1050 °C for 1 h. This demonstrates that the FCC phase is more stable than the BCC phase at low temperatures.
Acknowledgements The authors are pleased to acknowledge the financial support of this research by the National Science Council, People’s Republic of China under Grant NSC892216-E-011-024.
4. Conclusion The Fe-23.0 wt.% Mn-7.4 wt.% Al-0.03 wt.% C alloy has been intensively studied. After water quenching following 1 h at 1050 °C in air, the alloy has a small amount of austenite distributed discontinuously along the grain boundaries of the ferrite phase. When the as-quenched alloy was aged at higher temperatures, the precipitation of the austenite phase from the ferrite matrix grew preferentially along grain boundaries or within the matrix. The precipitation of the austenite phase within the ferrite matrix preferred the form of Widmansta¨ tten side-plates. The orientation relationships between the FCC Widmansta¨ tten sideplate and the BCC matrix are: (111)FCC//(011)BCC, [101]FCC//[111]BCC,
Fig. 5. The light micrograph (LM) of the Fe – Mn– Al alloy after heated at 1050 °C for 1 h then cooled at a rate of 50 °C h − 1 to room temperature (A, austenite; F, ferrite).
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Fig. 6. (a) The 200 dark-field image of the D03 phase from ferrite matrix after program-controlled furnace cooling. (b) SADP taken from a ferrite region. The zone axis is [011]. (hkl, ferrite phase; hkl, D03 phase; T, transmitted beam).
References [1] [2] [3] [4] [5]
P.R.S. Jackson, G.R. Wallwork, Oxid. Metals 21 (1984) 135. J.G. Duh, C.J. Wang, J. Mater. Sci. 25 (1990) 2063. K. Sato, K. Tanaka, Y. Inoue, ISIJ Int. 29 (1989) 788. T. Maki, C.M. Wayman, Acta Metall. 25 (1977) 681. M. Umemoto, T. Hyodo, T. Maeda, I. Tamura, Acta Metall. 32 (1984) 1191. [6] P.J. Brofman, G.S. Ansell, G.J. Judd, Metall. Trans. A 13 (1982) 203. [7] A. Sato, M. Kato, Y. Sunaga, T. Miyazaki, T. Mori, Acta Metall. 28 (1980) 367. [8] X.M. Zhang, E. Gautier, A. Simon, Acta Metall. 37 (1989) 477.
[9] R.F. Mehl, C.S. Barrett, D.W. Smith, Trans. AIME 105 (1953) 215. [10] D.A. Porter, K.E. Eastering, Phase Transformations in Metals and Alloys, second ed., 1992 [11] K.H. Hwang, C.M. Wan, J.G. Byrne, Scripta Metall. 24 (1990) 979. [12] S.K. Chen, K.H. Hwang, C.M. Wan, J.G. Byrne, Scripta Metall. 24 (1990) 151. [13] K.H. Hwang, C.M. Wan, J.G. Byrne, Mater. Sci. Eng. A 132 (1991) 161. [14] W.B. Lee, F.R. Chen, S.K. Chen, G.B. Olson, C.M. Wan, Acta Metall. 43 (1995) 21. [15] H.Y. Chu, F.R. Chen, T.B. Wu, Scripta Metall. 33 (8) (1995) 1269.
The precipitation of FCC phase from BCC matrix in an Fe–Mn–Al alloy Wei-Chun Cheng *, Hsin-Yu Lin Department of Mechanical Engineering, National Taiwan Uni6ersity of Science and Technology, Taipei 106, Taiwan Received 10 April 2001; received in revised form 15 May 2001
Abstract The Fe-23.0 wt.% Mn-7.4 wt.% Al-0.03 wt.% C alloy, after water quenching following 1 h at 1050 °C in air, has a small amount of austenite distributed discontinuously along the grain boundaries of the ferrite phase. Aging the as-quenched alloy at higher temperatures caused precipitation of the austenite phase from the ferrite matrix preferentially along grain boundaries or within the matrix. The precipitation of the austenite phase within the ferrite matrix preferred the form of Widmansta¨tten side-plates. Electron diffraction studies verify that the austenite phase precipitates from the ferrite matrix. A Kurdjumov– Sachs orientation relationship holds between the BCC matrix and the FCC precipitate. The proportion of the austenite phase is larger for lower temperatures aging or furnace cooling than in the as-quenched condition. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Quenching; Aging; Precipitation; Widmansta¨tten side-plate; Kurdjumov– Sachs orientation relationship
1. Introduction In recent decades Fe – Mn – Al alloys have gained much attention as substitutes for some of the conventional Fe –Ni –Cr stainless steels. Results concerning corrosion resistance and high temperature oxidation resistance have been reported in several publications [1,2]. For the development of corrosion resistance and high temperature oxidation resistance, the phase equilibria of the Fe – Mn – Al ternary system have practical importance. According to Sato’s work [3], in the range of 20–30 wt.% Mn and 4 – 20 wt.% Al at temperature of 1050 °C, Fe –Mn – Al alloys are composed of ferrite and austenite phases, i.e. some of them are single phase, either ferrite or austenite, and others are duplex phase. In common ferrous alloys, depending on the cooling rates for the austenite to ferrite phase transformations, Widmansta¨tten side-plate, massive, and martensite phases have been observed within the original FCC matrix during the high temperature quenching process [4 –9]. Experimental measurements on Widmansta¨tten * Corresponding author. Tel.: + 886-2-2737-6241; fax: + 886-22737-6460. E-mail address: [email protected] (W.-C. Cheng).
side-plate ferrite phase in common ferrous alloys show that the orientation relationships close to the Kurdjumov –Sachs or Nishiyama –Wasserman types are usually found [10]. In contrast to the common ferrous alloys in the phase transformation during cooling from high temperatures, Fe –Mn –Al alloys have gained much attention for the formation of the FCC phase inside the BCC parent phase. It is well known that among Fe –Mn –Al alloys, Mn and Al are FCC and BCC formers, respectively. Higher Mn content causes the higher proportion of the FCC phase to emerge at low temperatures in contrast to the full BCC phase in common ferrous alloys. If Fe –Mn –Al alloys contain low concentration of Al and high concentration of Mn, the full austenite phase could be observed even at room temperature [3]. In this unusual situation, it may seem that BCC (d-Fe) to FCC (g-Fe) phase transformation at high temperatures has been shifted to lower temperatures. But in common ferrous alloys, this kind of phase transformation is inhibited or screened out by the formation of another BCC (a-Fe) at low temperatures. Fortunately, the addition Mn into the ferrous alloy expands the FCC phase to low temperatures and the BCC (d-Fe) to FCC (g-Fe) phase transformation at higher temperatures can be
0921-5093/02/$ - see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 1 ) 0 1 4 9 8 - 8
W.-C. Cheng, H.-Y. Lin / Materials Science and Engineering A323 (2002) 462–466
shifted to room temperature. For BCC to FCC phase transformations in Fe– Mn – Al alloys, Widmansta¨ tten side-plate [11], massive [12], and 18R type martensitic phases [13–15] have also been observed within the original BCC matrix as the samples are quenched rapidly from 1300 °C. As the FCC phase is stable at low temperatures for some of the Fe– Mn – Al alloys, the purpose of the present work is to study the precipitation of the FCC phase within BCC matrix at low temperatures.
2. Experimental procedures Slabs with composition of Fe-23.0 wt.% Mn-7.4 wt.% Al-0.03 wt.% C were initially prepared by air induction melting. 1008 plain carbon steel, electrolytic manganese, and high– purity aluminum were cast into 10-kg ingots. After being homogenized at 1200 °C for 4 h under a protective argon atmosphere, the ingots were hot forged and then cold rolled to a thickness of 2 mm. These were heated at 1050 °C for 1 h in air, waterquenched to room temperature, then aged at higher temperatures for the study. In addition to the above treatment, samples were also heated at 1050 °C for 1 h in air then furnace cooled at a rate of 50 °C/h to room temperature via program control for comparison. Samples in the above heat treatments were sectioned, mechanically polished and etched in a 10% nital solution for light microscopy (LM) observation. Thin foils for transmission electron microscopy (TEM) were obtained by a twin-jet polisher in a 10% HClO4 and 90% ethanol solution at 30 V and 0.1 A cm − 2, examined in a JEOL 2000 FXII STEM and JEOL JEM 2010 TEM operated at 200 kV. Some of the samples were also examined in a RIGAKU DMAX-B X-ray diffractometer with the maximum power of 12 kW.
Fig. 1. Light micrograph (LM) of the Fe–Mn–Al alloy after quenching into room temperature water from 1 h at 1050 °C (A, austenite; F, ferrite).
463
3. Results and discussion Fig. 1 is the light micrograph (LM) of the Fe– Mn – Al alloy after being quenched into room temperature water following 1 h at 1050 °C. According to LM observation, the alloy consists of two phases: the matrix phase and the discrete phase along the grain boundary of the matrix. In an X-ray study, only ferrite and austenite phases are detected and confirmed by TEM study. The matrix phase is ferrite and the discrete phase along the grain boundaries of the matrix is austenite. The result is quite consistent with Sato’s observation in the two-phase region of Fe –Mn –Al alloys [3]. No precipitates are observed in the as-quenched condition by X-ray method, light or electron microscopy in either phase. The evolution of the precipitates during the aging process at higher temperatures was studied by transmission electron microscopy. In the present study, we conducted aging processes for temperatures ranging from 350 to 700 °C. Fig. 2(a) is a TEM bright-field image of the alloy after being aged at 600 °C for 12 h. It reveals the formation of the Widmansta¨tten sideplate precipitates within the ferrite matrix. The Widmansta¨tten precipitates were identified as the austenite phase, and the selected area diffraction patterns taken in these precipitates are shown in Fig. 3. This result was also confirmed by X-ray study in that only ferrite and austenite phases were detectable. The precipitation of the austenite phase preferred the grain boundaries of the ferrite phase and within the matrix, as shown in Fig. 2(b), in which the sample was aged at 540 °C for 24 h. According to the above observations, the phase transformation of the Fe –Mn –Al alloy during aging is opposite to that of conventional steels in that the low-temperature stable crystal structure is FCC rather than BCC. It is well known that orientation relationships between ferrite and austenite always accompany phase transformations in steel systems, for example martensitic transformation in Fe –Ni –C [4–8], and the formation of Widmannsta¨tten side-plates and intragranular phases in plain carbon steels [9]. Such orientation relationships provide information in which helps to understand the mechanisms of the structural changes. In the present Fe –Mn –Al alloy, during the BCC to FCC phase transformation, we found that the FCC Widmansta¨tten side-plate phase also had the orientation relationships with the parent matrix. Fig. 4(a –b) is selected area diffraction patterns (SADPs) of the FCC phase superimposed on that of the matrix. As indicated in Fig. 4(a –b), the orientation relationship between the FCC phase and the BCC matrix can be derived as: (111)FCC//(110)BCC, [101]FCC//[111]BCC,
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Fig. 2. TEM bright-field images of the alloy after being aged at 600 °C for 12 h (a) and at 540 °C for 24 h (b), respectively. (F, ferrite; G.B., grain boundary).
Fig. 3. The selected area diffraction patterns (SADPs) taken in the precipitates shown in Fig. 2 (a). Zone axes are: (a) [100]; (b) [111]; and (c) [011], respectively (T, transmitted beam).
[110]FCC near [100]BCC, which are the well known Kurdjumov– Sachs (K–S) orientation relationships [13]. This result is quite similar to that in the BCC to FCC phase transformations in Fe –Mn –Al alloys, Widmansta¨ tten side-plate [11], massive [12], and 18R type martensitic phases [13] which also have the same K– S orientation relationships with the original BCC matrix as the samples are quenched rapidly from 1300 °C. Aging the alloy at temperatures below 430 °C, caused the precipitation of D03 phase by homogeneous nucleation and growth and was first observed in the ferrite grains. For longer aging times, some Widmansta¨ tten precipitates were also present in addition to the D03 phase within the ferrite matrix. But at this stage they were too small to be identified. For aging temperatures above 430 °C, no D03 precipitates could
be observed, but the Widmansta¨ tten precipitates were observed under all circumstances. The precipitation of the FCC phase from the BCC matrix was observed during aging after water-quenching the slabs from 1 h at 1050 °C. It is clear that the proportion of the FCC phase was larger at lower temperatures than that at 1050 °C. In order to study the near equilibrium condition for the co-existence of the FCC and BCC phases at lower temperatures, we performed an experiment in which the slabs were heated at 1050 °C for 1 h and then cooled at a rate of 50 °C h − 1 to room temperature. Fig. 5 is the light micrograph (LM) of the Fe–Mn –Al alloy after being heated at 1050 °C for 1 h and then cooled at a rate of 50 °C h − 1 to room temperature. According to the LM observation, the FCC and BCC phases in the alloy are quite different from the condition shown in Fig. 1, and the proportion of the FCC phase is larger than that in
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465
Fig. 4. SADPs of the FCC phase superimposed on that of the BCC matrix in Fig. 2(b) showing the K – S orientation relationships. (a) Zone axes are [101]FCC and [111]BCC. (b) [110]FCC zone axis is near [100]BCC zone axis. (hkl, FCC; hkl, BCC).
the as-quenched condition. This observation demonstrates that the FCC phase is more stable than the BCC phase at low temperatures. In TEM study, D03 phase has formed within the ferrite phase during program-controlled furnace cooling as shown in Fig. 6(a), which is the g = 200 dark-field image of the D03 phase. Fig. 6(b) is the selected area diffraction pattern taken from the mixed region of ferrite matrix and the D03 fine precipitates. The zone axis is [011]. In addition to the spots corresponding to the ferrite matrix, the diffraction pattern also consists of small superlattice spots of 111, 200 etc. caused by the presence of the precipitates. The formation of the D03 phase within the ferrite phase is quite consistent with the result derived from the aging processes below 430 °C.
[110]FCC near [100]BCC, which correspond to the K–S orientation relationships. It is clear that the proportion of the austenite phase in the two-phase region of the Fe–Mn –Al alloy is larger at lower temperatures during aging or furnace cooling than that for the as-quenched condition from 1050 °C for 1 h. This demonstrates that the FCC phase is more stable than the BCC phase at low temperatures.
Acknowledgements The authors are pleased to acknowledge the financial support of this research by the National Science Council, People’s Republic of China under Grant NSC892216-E-011-024.
4. Conclusion The Fe-23.0 wt.% Mn-7.4 wt.% Al-0.03 wt.% C alloy has been intensively studied. After water quenching following 1 h at 1050 °C in air, the alloy has a small amount of austenite distributed discontinuously along the grain boundaries of the ferrite phase. When the as-quenched alloy was aged at higher temperatures, the precipitation of the austenite phase from the ferrite matrix grew preferentially along grain boundaries or within the matrix. The precipitation of the austenite phase within the ferrite matrix preferred the form of Widmansta¨ tten side-plates. The orientation relationships between the FCC Widmansta¨ tten sideplate and the BCC matrix are: (111)FCC//(011)BCC, [101]FCC//[111]BCC,
Fig. 5. The light micrograph (LM) of the Fe – Mn– Al alloy after heated at 1050 °C for 1 h then cooled at a rate of 50 °C h − 1 to room temperature (A, austenite; F, ferrite).
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Fig. 6. (a) The 200 dark-field image of the D03 phase from ferrite matrix after program-controlled furnace cooling. (b) SADP taken from a ferrite region. The zone axis is [011]. (hkl, ferrite phase; hkl, D03 phase; T, transmitted beam).
References [1] [2] [3] [4] [5]
P.R.S. Jackson, G.R. Wallwork, Oxid. Metals 21 (1984) 135. J.G. Duh, C.J. Wang, J. Mater. Sci. 25 (1990) 2063. K. Sato, K. Tanaka, Y. Inoue, ISIJ Int. 29 (1989) 788. T. Maki, C.M. Wayman, Acta Metall. 25 (1977) 681. M. Umemoto, T. Hyodo, T. Maeda, I. Tamura, Acta Metall. 32 (1984) 1191. [6] P.J. Brofman, G.S. Ansell, G.J. Judd, Metall. Trans. A 13 (1982) 203. [7] A. Sato, M. Kato, Y. Sunaga, T. Miyazaki, T. Mori, Acta Metall. 28 (1980) 367. [8] X.M. Zhang, E. Gautier, A. Simon, Acta Metall. 37 (1989) 477.
[9] R.F. Mehl, C.S. Barrett, D.W. Smith, Trans. AIME 105 (1953) 215. [10] D.A. Porter, K.E. Eastering, Phase Transformations in Metals and Alloys, second ed., 1992 [11] K.H. Hwang, C.M. Wan, J.G. Byrne, Scripta Metall. 24 (1990) 979. [12] S.K. Chen, K.H. Hwang, C.M. Wan, J.G. Byrne, Scripta Metall. 24 (1990) 151. [13] K.H. Hwang, C.M. Wan, J.G. Byrne, Mater. Sci. Eng. A 132 (1991) 161. [14] W.B. Lee, F.R. Chen, S.K. Chen, G.B. Olson, C.M. Wan, Acta Metall. 43 (1995) 21. [15] H.Y. Chu, F.R. Chen, T.B. Wu, Scripta Metall. 33 (8) (1995) 1269.