In the as-quenched condition, the microstructure of the alloy was (α＋γ) dual phases and no precipitates could be detected within the γ grains. Figure 5.1(a) is a selected-area diffraction pattern (SADP) taken from the α grain, exhibiting the presence of the superlattice reflection spots of the ordered D03
phase [19]. Figures 5.1(b) and (c) are (111) and (200) D03 dark-field (DF) electron micrographs, clearly showing the presence of extremely fine D03
domains and small B2 domains, respectively. The extremely fine D03 domains were formed within the α grain by a α → B2 → D03 continuous ordering transition during quenching [19]. This is similar to that observed by the present workers in the Fe-28.6wt.%Mn-10.1wt.%Al-0.46wt.%C alloy [10]. When the as-quenched alloy was aged at 650 for 6 h, fine k℃ ′ carbides were formed within the γ grains, as shown in Figure 5.2(a). Figure 5.2(b) is a bright-field (BF) electron micrograph taken from the α grain, indicating that some precipitates started to appear within the α grain. Figures 5.2 (c) and (d) are (111) and (200) D03 DF electron micrographs, indicating that the B2 domains grew significantly and the D03 domains remained to be extremely fine. Apparently, the B2 phase existed at the aging temperature and the extremely fine D03 domains were formed by a B2 → D03 ordering transition during quenching [19]. Figures 5.3(a)
5.2(b). According to the camera length and the measurement of angles as well as d-spacings of the diffraction spots, the crystal structure of the precipitate phase was determined to be monoclinic with lattice parameters a=1.158 nm, b=0.452 nm, c=0.509 nm, and β=98.3°, which corresponds to that of the Hägg carbide [18, 20-24]. It is clear that when the alloy was aged at 650 for 6 h,℃ the microstructure within the α grain was a mixture of (B2 + Hägg carbide). When the aging time was increased to 12 h, a new phase was found to occur on the α/α grain boundaries, as shown in Figure 5.4(a). Figure 5.4(b) is a SADP taken from the phase marked as “B” in Figure 5.4(a), indicating that the new phase on the α/α grain boundaries is β-Mn having a simple cubic structure with lattice parameter a=0.632 nm [25]. Figures 5.4(c) through (e) are three SADPs taken from the β-Mn marked as “B” and its surrounding B2 matrix, indicating that the orientation relationship between β-Mn and B2 phase was (100)B2 // (100)β-Mn, (010)B2 //(021)β-Mn, (011)B2 //(013)β-Mn, which is similar to that observed by the present workers in the Fe-29.9wt.%Mn-9.1wt.%Al-2.9wt.%Cr alloy [25]. When the aging time was increased to 18 h, the β-Mn would grow into adjacent B2 matrix, as illustrated in Figure 5.5(a). In Figure 5.5(a), it is also seen that another precipitate started to appear within the β-Mn region. Figure 5.5(b) is a BF electron micrograph taken from an area containing the precipitate and its surrounding β-Mn. Figures 5.5(c) through (e) are three SADPs taken from the
Figure 5.1 (a)
Figure 5.1 (b)
Figure 5.1 (c)
Figure 5.1 Transmission electron micrographs of the as-quenched alloy. (a) an SADP taken from the α grain. The zone axis is [011] (hkl: α phase, hkl: D03 phase). (b) and (c) (111) and (200) D03 DF images, respectively.
Figure 5.2 (a)
Figure 5.2 (b)
Figure 5.2 (c)
Figure 5.2 (d)
Figure 5.2 Transmission electron micrographs of the alloy aged at 650℃ for 6 h. (a) BF taken from the γ grain. (b) BF taken from the α grain. (c) and (d) (111) and (200) D03 DF images, respectively.
Figure 5.3 (a)
Figure 5.3 (b)
Figure 5.3 (c)
Figure 5.3 (a) through (c) Three SAPDs taken from the precipitate marked “C”
in Figure 5.2(b). The zone axes are (a) [110], (b) [101] and (c) ]
2 11
[ , respectively.
Figure 5.4 (a)
Figure 5.4 (b)
Figure 5.4 (c)
Figure 5.4 (d)
Figure 5.4 (e)
Figure 5.4 (a) SEM image of the alloy aged at 650℃ for 12 h. (b) an SADP taken from the precipitate marked “B” in Figure 5.4(a). The zone axis is [001]. (c) through (d) Three SADPs taken from an area including the precipitate marked “B” in Figure 5.4(a) and its surrounding matrix. The zone axes are (c) [001]β-Mn, [012]B2, (d) [012]β-Mn, [001]B2 and (e) [031]β-Mn, [011]B2, respectively.
Figure 5.5 (a)
Figure 5.5 (b)
Figure 5.5 (c)
Figure 5.5 (d)
Figure 5.5 (e)
Figure 5.5 (a) SEM image of the alloy aged at 650℃ for 18 h. (b) through (e) Transmission electron micrographs of the alloy aged at 650℃ for 18 h. (b) BF. (c) through (e) Three SADPs taken from the precipitate marked “P” in Figure 5.5(b). The zone axes are (c) [010], (d) [011] and (e) [111], respectively.
precipitate marked as “P” in Figure 5.5(b), indicating that the precipitate was M23C6-type carbide having a face-centered cubic (f.c.c) structure with lattice parameter a=1.079 nm [26]. After further prolonged aging, the β-Mn occurred to form inside the B2 matrix with different variants, as shown in Figure 5.6.
Based on the preceding results, some discussion is appropriate. That the phase transition (B2 + Hägg carbide) → (B2 + Hägg carbide + β-Mn) → (B2 + Hägg carbide + β-Mn + M23C6 carbide) occurred in the alloy aged at 650 is a ℃ remarkable feature in the present study. In order to clarify this feature, an TEM-EDS study was made. The chemical compositions of the Hägg carbide, M23C6 carbide, β-Mn and B2 phase in the alloy aged at 650 for differ℃ ent times are listed in Table 5.1. In the Fe-Al phase diagram [19], it is seen that when the Fe-Al alloy with Al 25.1 at.% is heat≧ -treated at 650 and then quenched to ℃ room-temperature, the microstructure present at 650 was B2 phase and a B2 ℃
→ D03 ordering transition would occur during quenching. The Al content of the B2 domain after being aged at 650 for 6 h is 25.28 at.%, as seen in Table ℃ 5.1.
Therefore, it is reasonable to suggest that the microstructure of the matrix at 650 would be the B2 phase and ℃ extremely fine D03 domains were formed by a B2 → D03 ordering transition during quenching, which is similar to that seen in
of the B2 domain in the alloy aged at 650 for 6 h℃ are much lower than those of the α phase in the as-quenched alloy, suggesting that the surrounding region would be enriched in Mn and Cr due to the growth of the B2 domains. It is reasonable to believe that the enrichment of Mn and Cr contents would enhance the formation of the (Mn,Cr)-rich Hägg carbide at the region contiguous to the B2 domains. Consequently, when the as-quenched alloy was aged at 650 for 6 ℃ h, the microstructure was a mixture of (B2 + Hägg carbide). After prolonged the aging time to 12 h, the β-Mn preferred to form heterogeneously on the α/α grain boundaries, which is similar to that observed in the FeMnAlCr alloys in the previous study [25]. In the Fe-Mn binary alloy, the β-Mn could only be found at temperature above 707 with Mn 70 at.%℃ ≧ [27]. However, the addition of Al would cause the β-Mn region to extend considerably [25]. Additionally, due to the formation of Mn-rich β-Mn, the (Mn,Cr)-rich Hägg carbide was observed to be disappeared within the β-Mn phase. Therefore, the microstructure of the aged at 650 for 12 h was a mixture of (B2 + ℃ Hägg carbide + β-Mn). Furthermore, when the aging time was increased to 18 h, the M23C6 carbides started to form within the β-Mn phase. Since the β-Mn has a simple cubic structure, the solubility of C could be expected to be low. Besides, it is well known that the Cr atom is a strong carbide former. In Table 5.1, it is clearly seen that the Cr content in the B2 phase in the alloy aged for 18 h were slightly less than that in
Figure 5.6
Figure 5.6 SEM image of the alloy aged at 650℃ for 36 h.
Table 5.1 Chemical compositions of the phases revealed by energy-dispersive
the alloy aged for 12 h. Therefore, the reason why M23C6 carbide occurred to form within the β-Mn phase is plausible that the C content was expected to be supersaturated and the Cr atom preferred to form a carbide phase as M23C6
rather than fully dissolve in the β-Mn phase. Consequently, the microstructure of the alloy aged at 650 for longer than 18 h was a mixture of (B2 + ℃ Hägg carbide + β-Mn + M23C6 carbide).
Finally, it is worthwhile pointing out that the presence of the M23C6 carbides within the β-Mn phase has never been observed by other workers before. The orientation relationship between the M23C6 carbide and the β-Mn phase has also not been reported in the FeMnAlC or FeMnAlCrC alloy systems in the previous literatures. Obviously, in order to clarify the orientation relationship between the M23C6 carbide and β-Mn phase, a further work is needed.
5-4 Conclusions
The as-quenched microstructure of the Fe-30wt.%Mn-10wt.%Al-4wt.%Cr- 0.45wt.%C alloy was (α＋γ) dual phases. No precipitates could be detected within the γ grains, whereas extremely fine D03 domains could be observed within the α grains. The extremely fine D03 domains were formed by a α → B2
→ D03 continuous ordering transition during quenching. When the as-quenched alloy was aged at 650 , fine ℃ κ′ carbides were precipitated within the γ grains.
Additionally, with increasing the aging time at 650 , a (B2 + ℃ Hägg carbide)→(B2 + Hägg carbide + β-Mn) → (B2 + Hägg carbide + β-Mn + M23C6
carbide) phase transition had occurred within the α grains. This phase transition has never been observed by other workers in the FeMnAlC and FeMnAlCrC alloy systems before.
References
1. K. Sato, K. Tagawa, Y. Inoue, Metall. Trans. A 21A (1990) 5.
2. W.K. Choo, J.H. Kim, J.C. Yoon, Acta Mater. 45 No.12 (1997) 4877.
3. C.Y. Chao, C.N. Hwang, T.F. Liu, Scripta Metall. 28 (1993) 109.
4. G.S. Krivonogov, M.F. Alekseyenko, G.G. Solovyeva, Fiz. Metal.
Metalloved 39 No.4 (1975) 775.
5. C.N. Hwang, C.Y. Chao, T.F. Liu, Scripta Metall. 28 (1993) 263.
6. T.F. Liu, J.S. Chou, C.C. Wu, Metall. Trans. A 21A (1990) 1891.
7. S.C. Tjong, S.M. Zhu, Mater. Trans. 38 No.2 (1997) 112.
8. S.C. Tjong, N.J. Ho, Metallography 21 (1988) 199.
9. T.F. Liu, C.M. Wan, Strength Met. Alloys 1 (1986) 423.
1589.
17. C.M. Liu, H.C. Cheng, C.Y. Chao, K.L. Ou, J. Alloys and Comp. 488 (2009) 52.
18. Y.H. Tuan, C.G. Chao, T.F. Liu, Mater. Trans. 51 No.6 (2010) 1168.
19. S.M. Allen, J.W. Cahn, Acta Metall. 24 (1976) 425.
20. F. Ernst, Y. Cao, G.M. Michal, Acta Mater. 52 (2004) 1469.
21. H.I. Faraoun, Y.D. Zhang, C. Esling, H. Aourag, J. App. Phys. 99 (2006) 093508.
22. E. Bauer-Grosse, C. Frantz, G. Lecaer, J. Non-Crys. Solids 44 (1981) 277.
23. V.D. Blank, Y.L. Alshevskiy, A.I. Zaitsev, N.V. Kazennov, I.A. Perezhogin, B.A. Kulnitskiy, Scripta Mater. 55 (2006) 1035.
24. M. Audier, P. Bowen, W. Jones, J. Crys. Growth 64 (1983) 291.
25. T.F. Liu, C.C. Wu, Scripta 23 (1989) 1087.
26. T.F. Liu, S.W. Peng, Y.L. Lin C.C. Wu, Metall. Trans. A 21 (1990) 567.
27. V.G. Rivlin, Inter. Metals Review 28 No.6 (1983) 309.