In the as-quenched condition, the microstructure of the alloy was (γ+α) dual phases. No precipitates could be detected within the γ grains; however, extremely fine D03 domains could be observed within the α grains. The extremely fine D03 domains were formed within the α grains by a continuous ordering transition during quenching. This is similar to that observed by the present workers in the Fe-28.6wt.%Mn-10.1 wt.%Al-0.46 wt.%C alloy [10].
When the as-quenched alloy was aged at 550 for 12 h, fine k′ carbides were ℃ formed within the γ grain, as shown in Figure 4.1(a). Figure 4.2(a) is a bright-field (BF) electron micrograph taken from the α grain, indicating that a lot of small precipitates occurred within the α matrix. Figure 4.2(b) is a selected-area diffraction pattern (SADP) taken from the α matrix, exhibiting the presence of the superlattice reflection spots of the ordered D03 phase [10].
Figures 4.2(c) and (d) are the (111) and (200) D03 dark-field (DF) electron micrographs taken from the same area as Figure 4.2(a), clearly revealing that the (111) DF image and the (200) DF image are morphologically identical.
Therefore, the bright domains presented in Figures 4.2(c) and (d) are of the D03
phase with a/2＜100＞ anti-phase boundaries (APBs) [10]. Figure 4.2(e) is the DF image of the small precipitates, indicating that these precipitates have occurred preferentially at a/2＜100＞ APBs of the D03 domains. A preliminary
Figure 4.1 BF transmission electron micrograph taken from the γ matrix of the alloy aged at 550 for 12 h.℃
Figure 4.2 (a)
Figure 4.2 (b)
Figure 4.2 (c)
Figure 4.2 (d)
Figure 4.2 (e)
Figure 4.2 Transmission electron micrographs taken from the D03 matrix of the alloy aged at 550 for 12 h. (a) BF, (b) an SADP, the zone axis ℃ is [011], (c) and (d) (111) and (200) D03 DF image, respectively, (e) (001)M5C2 DF image.
study of electron diffraction indicated that the precipitate was not of any known phases reported in FeMnAl, FeMnAlC and FeMnAlCrC alloy systems [1-17]. In order to clarify the crystal structure of the precipitate, eight SADPs taken from the precipitate marked as C in Figure 4.2(a) were obtained by tilting the specimen about some specific reflections. The results are shown in Figures 4.3(a)~(h). Table 4.1 shows the interplanar spacings of the precipitate phase, which were measured directly from the SADPs in Figures 4.3(a)~(h). The measured angles among the reciprocal lattice vectors are listed in Table 4.2.
Using these measured values of the interplanar spacings and angles, the crystal structure of the precipitate was determined to be monoclinic with lattice parameters a=1.158nm, b=0.452nm, c=0.509nm, and β=98.3°. Based on these lattice parameters, the interplanar spacings and the angles between the chosen reciprocal reflections were calculated by using the following equations [18]:
β
The calculated interplanar spacings and angles are also listed in Table 4.1
measured values are quite consistent with those obtained by calculation.
Therefore, the SADPs of the precipitate phase in Figures 4.3(a)~(h) could all be indexed. The zone axes of Figures 4.3(a)~(h) are [110], [210], [110], [312],
] 1
[10 , [314], [112] and [223], respectively. Compared with previous studies [19-23], it is clear that the crystal structure of the precipitate corresponds to that of the Hägg carbide (M5C2-type carbide).
Based on the preceding observations, two important experimental results are discussed below. ( ) The coarse MnⅠ -rich κ carbides or Mn-rich β-Mn precipitates were always observed within the α or D03 matrix in the duplex FeMnAlC alloys aged at 450~750 [8℃ -10]. However, only M5C2 carbides were formed at a/2＜100＞ APBs of the well-grown D03 domains, and no evidence of κ carbide or β-Mn precipitate could be detected within the α grain in the present alloy aged at 550 . In order to clarify the apparent difference, ℃ an TEM-EDS study was made. The average concentrations of the alloying elements were obtained from at least ten different EDS profiles of each phase. The results are summarized in Table 4.3. It is seen in Table 4.3 that the Cr and Mn concentrations in the well-grown D03 domains were much lower than those in the as-quenched alloy, and the reverse result was obtained in the M5C2 carbide.
Therefore, it is believed that during the growth of the D03 domains, partial Cr and Mn atoms would proceed to diffuse toward the a/2＜100＞ APBs. The
Figure 4.3 Eight SADPs taken from the precipitate marked “C” in Figure 4.2(a). The zone
axes are (a) [110], (b) [210], (c) [110], (d) [312], (e) [101], (f) [314], (g) [112] and
(h) [223], respectively.
Table 4.1 The d-spacings of the Hägg carbide.
Observed d-spacing＊ Calculated d-spacing＊＊ Indexed Plane
1 0.504 0.504 001
＊The observed d-spacings were measured from SADPs,
＊＊The calculated d-spacings were obtained on the basis of the monoclinic structure with lattice parameters a=1.158nm, b=0.452nm, c=0.509nm and
Table 4.2 Angles among some reciprocal vectors of the Hägg carbide.
Two Desired
Reciprocal Vectors Observed Angle＊ Calculated Angle＊＊
Fig. 3(a) (001) and (111) 51.8 51.9
Table 4.3 Chemical compositions of the phases revealed by EDS.
Chemical Composition (at.%) Heat
Treatment
Phase
Fe Mn Al Cr
γ 49.07 30.20 17.05 3.68
SHT
α + D03 52.07 21.80 21.55 4.58
D03 57.56 16.32 25.18 0.94
550 , 12h℃
M5C2 30.56 43.20 1.02 25.22
higher concentrations of both Cr and Mn would cause the (Cr,Mn)-rich M5C2
carbides to precipitate at a/2 ＜ 100 ＞ APBs. The precipitation of the (Cr,Mn)-rich M5C2 carbides would decrease the Mn concentration drastically, thus inhibiting the precipitation of both Mn-rich κ carbides and Mn-rich β-Mn precipitates within the α grain. ( II ) The Hägg carbide was extensively observed by many workers in the bcc-type alloys [20-24]. Depending on the chemical compositions, the lattice parameters of the Hägg carbide varied in the range of a=1.150-1.158nm, b=0.452-0.457nm, c=0.501-0.509nm and β=97.6-98.3°
[19-22]. However, to date, the orientation relationship between the Hägg carbide and the bcc-type (i.e. α, D03, B2) structure is very deficient. We are aware of two articles [22, 23], in which they reported that both Fe5C2 carbide and Fe3C carbide were formed intimately in α-iron after being heat-treated at 500-800 ℃ under a CO or/and H2 atmosphere. By using electron diffraction, the orientation relationship between the Fe5C2 and Fe3C was determined to be (100)Fe5C2 //
(001)Fe3C [22, 23]. In addition, they correlated the obtained result with the orientation relationship between Fe3C and α phase, (001)Fe3C // (211)α, which was reported by other workers in ferritic stainless steel [24]. Therefore, they deduced that the orientation relationship among Fe5C2, Fe3C and α phase was (100)Fe5C2 //
(001)Fe3C // (211)α [23]. It is well-known that the orientation relationship between
parallel planes or two pairs of parallel planes. However, in the previous studies [22 ,23], only a pair of parallel planes was determined and no direct experiment evidence confirmed the orientation relationship between Fe5C2 and α phase.
Therefore, the electron diffraction technique was used to clarify the orientation relationship between the M5C2 carbide and the D03 matrix in the present study.
The results are presented in Figures 4.4(a) and (b). In these two SADPs, it is clear that the (510) and (134) reflection spots of the M5C2 carbide are parallel to the (110) and (102) reflection spots of the D03 matrix, respectively. Accordingly, the orientation relationship between the M5C2 carbide and D03 matrix can be stated as follows: (510)M5C2 //(110)D03 and (134)M5C2
//(102)D03. In order to further certify the determined orientation relationship, the angle between the (510)M5C2 and (134) M5C2 was calculated by using the equation mentioned above. The calculated angle was 71.61°, which is quite consistent with the angle of 71.57° between the (110)D03 and (102)D03. Finally, it is worth mentioning that in the present study, a lot of effort was made to determine the parallel relationship of lower index planes between the M5C2
carbide and D03 matrix. Unfortunately, the attempt was not successful.
Figure 4.4 (a)
Figure 4.4 (b)
Figure 4.4 Two SADPs taken from an area including the precipitate marked
“C” in Figure 4.2(a) and its surrounding matrix. The zone axes are (a) [152]M5C2, [110]D03 and (b) [154]M5C2, [221] D03, respectively.
4-4 Conclusions
In Summary, the as-quenched microstructure of the Fe-30wt.%Mn- 10wt.%Al-4wt.%Cr-0.45wt.%C alloy was (γ+α) dual phases, and extremely fine D03 domains could be observed within the α grains. After being aged at 550 ℃ for 12 h, fine κ′ carbides were formed within the γ grains and the (Cr,Mn)-rich Hägg carbides occurred at a/2＜100＞ APBs of the well-grown D03 domains.
The Hägg carbide has never been observed by previous workers in FeMnAlC and FeMnAlCrC alloy systems. The orientation relationship between the Hägg carbide and D03 matrix was determined to be (510)M5C2 //(110)D03 and
) 4 13
( M5C2 //(102)D03. The orientation relationship between Hägg carbide and bcc-type phase has also never been reported by other workers before.
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