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Dynamic recrystallization of Nb3Al produced from alloy powder
Materials Science and Engineering, A159 (1992) 173-180
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Dynamic recrystallization of Nb3A1 produced from alloy powder Y. Murayama, S. Hanada and K. Obara Institute for Materials Research, Tohoku University, Sendai (Japan) (Received May 15, 1992)
Abstract The deformation behaviour of Nb3AI compressed at above 1300 K was studied using off-stoichiometric, single-phase Nb3A1 produced by consolidation of prealloyed and ball-milled powders. It was found that Nb3AI was able to be deformed plastically above 1400 K and the stress-strain curves showed deformation softening after a peak stress. Transmission electron microscopic observations revealed that this deformation softening was caused by dynamic recrystallization. The amount of deformation softening decreases and steady state deformation occurs at an early stage of compression, as the compression temperature increases.
1. Introduction The intermetallic compound Nb3AI has attracted attention for high temperature structural applications, because Nb3Al-based alloys have high strength and high creep resistance at elevated temperatures [1]. However, very little information is available on the mechanical properties of stoichiometric, single-phase Nb3AI, since it is difficult to prepare microcrack-free specimens by ingot metallurgy. Therefore, we have studied the high temperature deformation of offstoichiometric Nb3AI produced by the consolidation of prealloyed and ball-milled powders [2, 3]. Active slip systems in Nb3AI and Nb3Al-based alloys are 1/2(001){100} or (001){100}. Therefore, ductile plastic deformation of a Nb3AI polycrystal cannot be attained only by slip deformation, since the operation of these slip systems does not satisfy the von Mises criteria. However, Murayama et al. [3] have recently revealed that single-phase Nb3AI deforms plastically to large strains in compression above about 1400 K and shows a deformation structure indicating dynamic recrystallization (DRX). The DRX has been observed in many intermetallic compounds [4-20]. Hanada et al. [5, 6] pointed out the important role of DRX in the superplasticity of Fe3(A1,Si), the intermetallic compounds in which they first found superplasticity. Superplasticity has been found in some other intermetallic compounds, e.g. Ni3A1 [7, 8, 10], TiAI [15, 16] and Ni3Si [18, 19]. However, no decisive mechanism to explain the superplasticity has been presented, although various accommodation processes at grain boundaries have been discussed. This work is 0921-5093/92/$5.00
intended to study the deformation behaviour of Nb3AI compressed above 1300 K, using off-stoichiometric Nb3AI produced by the consolidation of prealloyed and ball-milled powders. In particular, we focus on the role of DRX for the ductilization of Nb3A1.
2. Experimental procedure Nb-AI alloy buttons with nominal stoichiometric composition were prepared by arc melting in an argon atmosphere. The buttons were crushed and milled to under 200 mesh size in air. The atomic ratio of niobium to aluminium in the alloy powders was 75 : 25 as determined by chemical analysis. The X-ray diffraction patterns of the alloy powders showed mainly strong peaks of Nb3A1 and a few weak peaks of Nb2AI. The powders were compacted at room temperature and then sintered at 1673, 1873 and 2073 K to determine an appropriate temperature for hot isostatic pressing (HIP) treatment. On increasing the sintering temperature of the compacts, the peaks of Nb2A1 disappeared, and only peaks of Nb3A1 were observed above 1873 K. Optical microscopy indicated that densification occurred to some degree during sintering at 1873 K. Therefore, 1873 K was selected as the temperature for HIP treatment. The green compacts were canned in evacuated glass tubes at 1673 K and then hot isostatically pressed at 1873 K and 147 MPa for 1.8 ks. The X-ray diffraction pattern of these specimens showed only peaks of Nb3AI. The atomic ratio of niobium to aluminium of the HIP-treated specimen was 76.8 : 23.2 as determined by chemical analysis. The HIP-treated © 1992 - Elsevier Sequoia. All rights reserved
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specimen consisted of off-stoichiometric Nb3AI (over 93 vol.%) and a small amount of niobium solid solution (3.7 vol.%) and AI203 (2.5 vol.%). Although niobium solid solution and AI203 were present, the concentration of aluminium in the Nb3AI matrix ranged from 23 to 26 at.% as determined by transmission electron microscopy with energy-dispersive X-ray spectroscopy (TEM-EDX). Therefore, the specimen used in this work is off-stoichiometric Nb3AI of almost single phase with a grain size of 11 ~m. Compression samples with dimensions 2 m m × 2 mm × 3.7 mm were prepared from the HIP-treated specimens. Compression tests were carried out above 1273K at an initial strain rate of 2.25×10 -4 s -~ under a vacuum of about 2 × 10 - 3 Pa. The samples were compressed between SiC plates with graphite for lubrication. The deformation mechanism was investigated by TEM. Details of the preparation of the thin foil for TEM are given elsewhere [2].
3. Results
Figure 1 shows stress-strain curves of HIP-treated specimens. The sample strained at 1300 K broke before yielding but showed a high fracture strength near 2GPa. At 1400 and 1490K, the samples deformed without cracking macroscopically at the early stages of deformation. At large strains, cracks parallel to the compression direction appeared in the barrel-shaped sample. Although the flow stress decreased significantly with increasing temperature, the high peak stress near 1 GPa was obtained even at 1490 K. At 1650 and 1800 K, there was steady state deformation. The samples deformed at 1490 and 1650 K were unloaded at the strains indicated by
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Fig. 1. Stress-strain curves of the HIP-treated specimens. Arrows indicate interrupted strains for observation of structure.
arrows to investigate the deformation microstructure by TEM. Some preliminary experiments showed that there were no dislocations in the samples before compression. At the strain showing the peak stress, dislocations with banded structure were arranged along the extended faults in the specimen compressed at 1490 K as shown in Fig. 2. Tangled dislocations with a high density can be seen in the band. Rectangular arrangements of dislocations stretched in the (100) directions were visible and became well defined with increasing strain. The (100){001} dislocation induced by the plastic deformation dissociates into two 1/2(100){001} partial dislocations with a complex stacking fault whose displacement vector is 1/2(100) [2]. At this strain, new grains could already be seen at grain boundaries. In the specimen compressed to the strain of 0.07 at 1490 K, where the flow stress decreased significantly, the interiors of the grains were covered with rectangularly arranged dislocations and new recrystallized grains developed along the initial grain boundaries (Fig. 3). The micrograph in Fig. 3 indicates that DRX occurs at or near the initial grain boundaries and a necklace structure is created in a similar manner to a lot of dynamically recrystallized metals. The diffraction patterns of deformed grains contain the streaks along the (100) direction owing to the stacking faults. The stacking faults and the partial dislocations also can be seen in the new grains. At about 0.33 strain, the deformation microstructure is mainly composed of fine recrystallized grains less than 1/~m in size (Fig. 4). In addition to the fine recrystallized grains, many stacking faults and dislocations can be seen in some grains. The grains containing a lot of faults are regarded as deformed initial grains, because the grain size is larger than that of the new grains. Although the rate of deformation softening decreased and the stress-strain curves became horizontal at large strains, none of the initial grains has been replaced by dynamically recrystallized grains. This suggests that grain-boundary sliding operates around grains produced by DRX. The stress-strain curve of the sample compressed at 1650 K shows steady state flow from an early stage o f the deformation. Figure 5 shows a micrograph of the sample compressed at 1650 K to the strain showing the peak stress. Dynamically recrystallized grains also can be seen in the sample. A large grain with many stacking faults and dislocations can be seen at the lower part of the micrograph. Figure 6 shows the dislocation structure in this grain. It is evident that many stacking faults are bounded by the partial dislocations and that subboundaries are formed by these faults. The widths of the stacking faults between the partial dislocations are narrower than those in the sample compressed at
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Fig. 2. Transmission electron micrograph of a sample compressed at 1490 K to a strain showing peak stress.
1490 K. In the sample strained to about 0.3, the deformation structure consists of dislocation-free dynamically recrystallized grains and subgrains. Two kinds of subboundaries can be seen; the first is composed of arrays of short partial dislocations (as shown in Fig. 6) and the second is made of long straight dislocations (as shown in Fig. 7).
4. Discussion
4.1. Tefnperature dependence of D R X Nb3AI produced by powder metallurgy in this experiment deformed plastically above 1400 K at an initial strain rate of 2.25 x 10 -4 s -1. It is recognized that this plastic deformation at the early stage is caused
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Fig. 3. Transmissionelectron micrographof a sample compressedat 1490 K to a strain showingdeformation softening.
by 1/2(001){100} partial dislocations with the complex stacking fault of displacement vector 1/2(001). Although piling-up of dislocations at a grain boundary enables dislocations to be generated in an adjacent grain (Fig. 8), the adjacent grain cannot accommodate large strain at the grain boundary because of a deficient slip system. That is, plastic deformation of an Nb3AI polycrystal cannot be maintained by these slip systems
alone. Unless some accommodation process acts at the grain boundaries, intergranular fracture or cleavage fracture must be induced by stress concentration. Thus, DRX occurs as an accommodation process in this experiment. DRX in metals and alloys has been extensively studied [21]. When metals deform plastically by crystallographic slip at elevated temperatures, restoration
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Fig. 4. Transmissionelectron micrograph of a sample compressedat 1490 K to a large strain of 0.33.
processes and strain hardening processes operate simultaneously. Dynamic recovery (DRC) and DRX are regarded as the restoration processes and are exclusive to each other. Hardwick et al. proposed that the priority of DRC and DRX depends on the magnitude of the stacking fault energy [22]. In metals with high stacking fault energies, e.g. A1, the flow stress is saturated to a steady state condition owing to the balance between DRC and strain hardening without DRX, because dislocations are easily rearranged by dislocation climb and form subgrains. However, the priority of DRC and DRX depends not only on the stacking fault energy but also on the deformation condition [23, 24]. McQueen [25] has recently claimed that microstructure should be considered as the criterion for DRX. His idea is concerned with the
development of a dislocation wall of a boundary which is capable of migration and is associated with a small region of large misorientation (more than 10"). Furthermore, he stresses that there is a sufficient difference in stored energy between that in the volume within the nucleus and that in the surroundings. Equivalent 1/2(001){ 100} dislocations in the sample compressed at 1490 K react with each other and form rectangular networks along the direction of the planar faults during the early stage of deformation in our experiment (Fig. 2). As the deformation proceeds, these rectangular networks completely cover the interior of all the grains. Therefore, the misorientation at the wall of the rectangular networks seems to develop by dislocation climbing. Since the stored energy near a grain boundary is higher than that inside
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Fig. 5. Transmission electron micrograph of a sample compressed at 1650 K to a strain showing peak stress.
Fig. 6. Dislocation structure of an initial grain shown in the micrograph of Fig. 5.
Fig. 7. Transmission electron micrograph of a sample compressed at 1650 K to a strain of 0.31.
a grain, it is expected that the wall of networks near a grain boundary migrates and new recrystallized grains are produced. As shown in Fig. 2, the width of the complex stacking faults is very large because of the low stacking fault energy in the sample of our experiment. This situation is disadvantageous for D R C which requires dislocation climbing, and these conditions suppress D R C and
induce DRX. It should be noted that the formation of networks and the misorientation development between those cells play important roles in DRX. In the sample compressed at 1650 K, rectangular networks cannot be seen clearly. The width of the complex stacking faults is narrower than that at 1490 K and subboundaries were observed as shown in Figs. 6 and 7. This means that the stacking fault energy
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phase in a two-phase stainless steel refined the grain size of the a phase and then the resultant grain boundary sliding caused superplasticity. On the other hand, Takasugi et al. [18, 19] indicated that grain boundary sliding is a main deformation mechanism in superplastic Ni3(Si,Ti ). In Zr3AI [4] and (Fe, Co)3V [12], the grain boundary sliding of DRX grains was found to lead to ductilization. In these experiments, grain boundary sliding was confirmed by the large misorientation between adjacent grains and the existence of dislocation-free grains. The mixed structure of D R X grains with and without dislocations can be seen in the sample compressed at 1490 K to the strain of 0.33 in our experiment. In addition, there exist large grains with a high dislocation density which are considered to be initial grains. This indicates that none of the initial grains has been replaced by D R X grains, although D R X grains are predominant. Thus, it is concluded that the ductile deformation in Nb3AI polycrystal is due to the occurrence of DRX at initial grain boundaries and subsequent grain boundary sliding.
5. Summarizing remarks
Fig. 8. Dislocations near a grain boundary of a sample compressed at 1490 K to a strain of 0.07. increases with increasing temperature. New grains formed by DRX were observed less frequently at grain boundaries in the sample compressed at 1650 K than in that compressed at 1490 K. These microstructural observa.tions are consistent with the fact that, as the compression temperature increases, the amount of deformation softening decreases. The sample compressed at 1650 and 1800 K did not exhibit a clear peak stress. The flow stress shifts immediately from the stage of initial strain hardening to steady state deformation. This is because the restoration process occurs rapidly at these temperatures. 4.2. Role o f D R X f o r ductifity
Since superplasticity was found in Fe3(AI , Si) [5, 6], superplasticity has been reported in many intermetallic compounds [10, 15, 16, 18-20, 26-28]. Superplasticity in most intermetallic compounds has been discussed in terms of DRX. Recently, D R X has been found to play an important role in the superplasticity of some disordered alloys. Maehara [29] proposed that superplasticity is caused by the balance between D R X and the strain hardening in a duplex phase stainless steel. Tsuzaki et al. [30] suggested that D R X of the a
Off-stoichiometric, single-phase Nb3AI produced by powder metallurgy deformed plastically above 1400 K and a ductile deformation behaviour without macroscopic cracking appeared above 1490 K. The stress-strain curves of the sample showed deformation softening owing to DRX. It is revealed that the formation of complex stacking faults bounded by partial dislocations and characteristic dislocation networks and subboundaries results in DRX.
References 1 D. L. Anton, D. M. Shah, D. N. Duhl and A. E Giamei, J. Organomet. Chem., 41 (2)(1989) 12. 2 Y. Murayama, S. Hanada, K. Obara and K. Hiraga, Philos. Mag. A , in press. 3 Y. Murayama, S. Hanada, K. Obara, K. Hiraga and S. Ishida, in O. Izumi (ed.), Proc. Int. Symp. on Intermetallic Compounds, 1991, Jpn. Inst. Met., Sendai, p. 1033. 4 E.M. Schulson, Metall. Trans. A, 9(1978) 527. 5 S. Hanada, T. Sato, S. Watanabe and O. Izumi, J. Jpn Inst. Met., 45(1981) 1285. 6 S. Hanada, T. Sato, S. Watanabe and O. Izumi, J. Jpn Inst. Met., 45(1981) 1293. 7 M.S. Kim, S. Hanada, S. Watanabe and O. Izumi, J. Jpn Inst. Met., 52 (1988) 1020. 8 M. S. Kim, S. Hanada, S. Watanabe and O. Izumi, Metall. Trans. JIM, 30 (1989) 77. 9 I. Baker, D. V. Viens, E. M. Schulson, Scr. Metall., 25 (1991) 889.
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10 A. Choudhury, A. K. Mukherjee and V. K. Sikka, J. Mater. Sci., 25 (1990) 3142. 11 G. Gottstein, L Chang and H. F. Yung, Mater Sci. TechnoL, 7(1991) 158. 12 C. T. Liu and E. M. Schulson, Metall. Tram. A, 15 (1984) 701. 13 M. Nobuki, K. Hashimoto, T. Tsujimoto and Y. Asai, J. Jpn Inst. Met., 50 (1986) 840. 14 H. Fukutomi, C. Hartig and H. Mecking, Z. Metallkd., 81 (1990) 272. 15 G.A. Salishchev, R. M. Galeyev and R. M. Imaev, in S. Hori, M. Tokizane and N. Furushiro (eds.), Superplasticity in Advanced Materials, Jpn. Soc. Res. Superplasticity, Osaka, 1991, p. 163. 16 K. Ameyama, A. Miyazaki and M. Tokizane, in S. Hori, M. Tokizane and N. Furushiro (eds.), Superplasticity in Materials, Jpn. Soc. Res. Superplasticity, Osaka, 1991, p. 317. 17 Y. Liu, T. Takasugi, O. Izumi and H. Suenaga, J. Mater. Sci., 24(1989) 4458. 18 T. Takasugi, S. Rikukawa and S. Hanada, Scr. Metall., 25 (1991)889.
19 T. Takasugi, S. Rikukawa and S. Hanada, in S. Hori, M. Tokizane and N. Furushiro (eds.), Superplasticity in Advanced Materials, Jpn. Soc. Res. Superplasticity, Osaka, 1991, p. 329. 20 S. Hanada, in S. Hod, M. Tokizane and N. Furushiro (eds.), Superplasticity in Advanced Materials, Jpn. Soc. Res. Superplasticity, Osaka, 1991, p. 289. 21 C. M. Sellars, Philos. Trans. R. Soc. London, Ser. A, 288 (1978) 147. 22 D. Hardwick and W. J. McG. Tegart, J. Inst. Met., 90 (1961) 17. 23 G. Gottstein and S. Deshpande, Mater. Sci. Eng., 94 (1987) 147. 24 G. Glover and C. M. Sellars, Metall. Trans.,4(1973)765. 25 H.J. McQueen, Mater. Sci. Eng., AIOI (1988) 149. 26 R.N. Wright and V. K. Sikka, J. Mater. Sci., 23 (1988) 4315. 27 T.G. Nieh andW. C. Oliver, Scr. Metall., 23(1989) 851. 28 J. Mukhopadhyay, G. Kasclmer and A. K. Mukheoee, Scr. Metall., 24 (1990) 857. 29 Y. Maehara, Tram. Iron Steellmt. Jpn, 27(1987) 705. 30 K. Tsuzaki, H. Matsuyama, M. Nagao and T. Maki, J. Jpn Inst. Met., 54 (1990) 878.
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Dynamic recrystallization of Nb3A1 produced from alloy powder Y. Murayama, S. Hanada and K. Obara Institute for Materials Research, Tohoku University, Sendai (Japan) (Received May 15, 1992)
Abstract The deformation behaviour of Nb3AI compressed at above 1300 K was studied using off-stoichiometric, single-phase Nb3A1 produced by consolidation of prealloyed and ball-milled powders. It was found that Nb3AI was able to be deformed plastically above 1400 K and the stress-strain curves showed deformation softening after a peak stress. Transmission electron microscopic observations revealed that this deformation softening was caused by dynamic recrystallization. The amount of deformation softening decreases and steady state deformation occurs at an early stage of compression, as the compression temperature increases.
1. Introduction The intermetallic compound Nb3AI has attracted attention for high temperature structural applications, because Nb3Al-based alloys have high strength and high creep resistance at elevated temperatures [1]. However, very little information is available on the mechanical properties of stoichiometric, single-phase Nb3AI, since it is difficult to prepare microcrack-free specimens by ingot metallurgy. Therefore, we have studied the high temperature deformation of offstoichiometric Nb3AI produced by the consolidation of prealloyed and ball-milled powders [2, 3]. Active slip systems in Nb3AI and Nb3Al-based alloys are 1/2(001){100} or (001){100}. Therefore, ductile plastic deformation of a Nb3AI polycrystal cannot be attained only by slip deformation, since the operation of these slip systems does not satisfy the von Mises criteria. However, Murayama et al. [3] have recently revealed that single-phase Nb3AI deforms plastically to large strains in compression above about 1400 K and shows a deformation structure indicating dynamic recrystallization (DRX). The DRX has been observed in many intermetallic compounds [4-20]. Hanada et al. [5, 6] pointed out the important role of DRX in the superplasticity of Fe3(A1,Si), the intermetallic compounds in which they first found superplasticity. Superplasticity has been found in some other intermetallic compounds, e.g. Ni3A1 [7, 8, 10], TiAI [15, 16] and Ni3Si [18, 19]. However, no decisive mechanism to explain the superplasticity has been presented, although various accommodation processes at grain boundaries have been discussed. This work is 0921-5093/92/$5.00
intended to study the deformation behaviour of Nb3AI compressed above 1300 K, using off-stoichiometric Nb3AI produced by the consolidation of prealloyed and ball-milled powders. In particular, we focus on the role of DRX for the ductilization of Nb3A1.
2. Experimental procedure Nb-AI alloy buttons with nominal stoichiometric composition were prepared by arc melting in an argon atmosphere. The buttons were crushed and milled to under 200 mesh size in air. The atomic ratio of niobium to aluminium in the alloy powders was 75 : 25 as determined by chemical analysis. The X-ray diffraction patterns of the alloy powders showed mainly strong peaks of Nb3A1 and a few weak peaks of Nb2AI. The powders were compacted at room temperature and then sintered at 1673, 1873 and 2073 K to determine an appropriate temperature for hot isostatic pressing (HIP) treatment. On increasing the sintering temperature of the compacts, the peaks of Nb2A1 disappeared, and only peaks of Nb3A1 were observed above 1873 K. Optical microscopy indicated that densification occurred to some degree during sintering at 1873 K. Therefore, 1873 K was selected as the temperature for HIP treatment. The green compacts were canned in evacuated glass tubes at 1673 K and then hot isostatically pressed at 1873 K and 147 MPa for 1.8 ks. The X-ray diffraction pattern of these specimens showed only peaks of Nb3AI. The atomic ratio of niobium to aluminium of the HIP-treated specimen was 76.8 : 23.2 as determined by chemical analysis. The HIP-treated © 1992 - Elsevier Sequoia. All rights reserved
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specimen consisted of off-stoichiometric Nb3AI (over 93 vol.%) and a small amount of niobium solid solution (3.7 vol.%) and AI203 (2.5 vol.%). Although niobium solid solution and AI203 were present, the concentration of aluminium in the Nb3AI matrix ranged from 23 to 26 at.% as determined by transmission electron microscopy with energy-dispersive X-ray spectroscopy (TEM-EDX). Therefore, the specimen used in this work is off-stoichiometric Nb3AI of almost single phase with a grain size of 11 ~m. Compression samples with dimensions 2 m m × 2 mm × 3.7 mm were prepared from the HIP-treated specimens. Compression tests were carried out above 1273K at an initial strain rate of 2.25×10 -4 s -~ under a vacuum of about 2 × 10 - 3 Pa. The samples were compressed between SiC plates with graphite for lubrication. The deformation mechanism was investigated by TEM. Details of the preparation of the thin foil for TEM are given elsewhere [2].
3. Results
Figure 1 shows stress-strain curves of HIP-treated specimens. The sample strained at 1300 K broke before yielding but showed a high fracture strength near 2GPa. At 1400 and 1490K, the samples deformed without cracking macroscopically at the early stages of deformation. At large strains, cracks parallel to the compression direction appeared in the barrel-shaped sample. Although the flow stress decreased significantly with increasing temperature, the high peak stress near 1 GPa was obtained even at 1490 K. At 1650 and 1800 K, there was steady state deformation. The samples deformed at 1490 and 1650 K were unloaded at the strains indicated by
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Fig. 1. Stress-strain curves of the HIP-treated specimens. Arrows indicate interrupted strains for observation of structure.
arrows to investigate the deformation microstructure by TEM. Some preliminary experiments showed that there were no dislocations in the samples before compression. At the strain showing the peak stress, dislocations with banded structure were arranged along the extended faults in the specimen compressed at 1490 K as shown in Fig. 2. Tangled dislocations with a high density can be seen in the band. Rectangular arrangements of dislocations stretched in the (100) directions were visible and became well defined with increasing strain. The (100){001} dislocation induced by the plastic deformation dissociates into two 1/2(100){001} partial dislocations with a complex stacking fault whose displacement vector is 1/2(100) [2]. At this strain, new grains could already be seen at grain boundaries. In the specimen compressed to the strain of 0.07 at 1490 K, where the flow stress decreased significantly, the interiors of the grains were covered with rectangularly arranged dislocations and new recrystallized grains developed along the initial grain boundaries (Fig. 3). The micrograph in Fig. 3 indicates that DRX occurs at or near the initial grain boundaries and a necklace structure is created in a similar manner to a lot of dynamically recrystallized metals. The diffraction patterns of deformed grains contain the streaks along the (100) direction owing to the stacking faults. The stacking faults and the partial dislocations also can be seen in the new grains. At about 0.33 strain, the deformation microstructure is mainly composed of fine recrystallized grains less than 1/~m in size (Fig. 4). In addition to the fine recrystallized grains, many stacking faults and dislocations can be seen in some grains. The grains containing a lot of faults are regarded as deformed initial grains, because the grain size is larger than that of the new grains. Although the rate of deformation softening decreased and the stress-strain curves became horizontal at large strains, none of the initial grains has been replaced by dynamically recrystallized grains. This suggests that grain-boundary sliding operates around grains produced by DRX. The stress-strain curve of the sample compressed at 1650 K shows steady state flow from an early stage o f the deformation. Figure 5 shows a micrograph of the sample compressed at 1650 K to the strain showing the peak stress. Dynamically recrystallized grains also can be seen in the sample. A large grain with many stacking faults and dislocations can be seen at the lower part of the micrograph. Figure 6 shows the dislocation structure in this grain. It is evident that many stacking faults are bounded by the partial dislocations and that subboundaries are formed by these faults. The widths of the stacking faults between the partial dislocations are narrower than those in the sample compressed at
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Fig. 2. Transmission electron micrograph of a sample compressed at 1490 K to a strain showing peak stress.
1490 K. In the sample strained to about 0.3, the deformation structure consists of dislocation-free dynamically recrystallized grains and subgrains. Two kinds of subboundaries can be seen; the first is composed of arrays of short partial dislocations (as shown in Fig. 6) and the second is made of long straight dislocations (as shown in Fig. 7).
4. Discussion
4.1. Tefnperature dependence of D R X Nb3AI produced by powder metallurgy in this experiment deformed plastically above 1400 K at an initial strain rate of 2.25 x 10 -4 s -1. It is recognized that this plastic deformation at the early stage is caused
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Fig. 3. Transmissionelectron micrographof a sample compressedat 1490 K to a strain showingdeformation softening.
by 1/2(001){100} partial dislocations with the complex stacking fault of displacement vector 1/2(001). Although piling-up of dislocations at a grain boundary enables dislocations to be generated in an adjacent grain (Fig. 8), the adjacent grain cannot accommodate large strain at the grain boundary because of a deficient slip system. That is, plastic deformation of an Nb3AI polycrystal cannot be maintained by these slip systems
alone. Unless some accommodation process acts at the grain boundaries, intergranular fracture or cleavage fracture must be induced by stress concentration. Thus, DRX occurs as an accommodation process in this experiment. DRX in metals and alloys has been extensively studied [21]. When metals deform plastically by crystallographic slip at elevated temperatures, restoration
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Fig. 4. Transmissionelectron micrograph of a sample compressedat 1490 K to a large strain of 0.33.
processes and strain hardening processes operate simultaneously. Dynamic recovery (DRC) and DRX are regarded as the restoration processes and are exclusive to each other. Hardwick et al. proposed that the priority of DRC and DRX depends on the magnitude of the stacking fault energy [22]. In metals with high stacking fault energies, e.g. A1, the flow stress is saturated to a steady state condition owing to the balance between DRC and strain hardening without DRX, because dislocations are easily rearranged by dislocation climb and form subgrains. However, the priority of DRC and DRX depends not only on the stacking fault energy but also on the deformation condition [23, 24]. McQueen [25] has recently claimed that microstructure should be considered as the criterion for DRX. His idea is concerned with the
development of a dislocation wall of a boundary which is capable of migration and is associated with a small region of large misorientation (more than 10"). Furthermore, he stresses that there is a sufficient difference in stored energy between that in the volume within the nucleus and that in the surroundings. Equivalent 1/2(001){ 100} dislocations in the sample compressed at 1490 K react with each other and form rectangular networks along the direction of the planar faults during the early stage of deformation in our experiment (Fig. 2). As the deformation proceeds, these rectangular networks completely cover the interior of all the grains. Therefore, the misorientation at the wall of the rectangular networks seems to develop by dislocation climbing. Since the stored energy near a grain boundary is higher than that inside
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Fig. 5. Transmission electron micrograph of a sample compressed at 1650 K to a strain showing peak stress.
Fig. 6. Dislocation structure of an initial grain shown in the micrograph of Fig. 5.
Fig. 7. Transmission electron micrograph of a sample compressed at 1650 K to a strain of 0.31.
a grain, it is expected that the wall of networks near a grain boundary migrates and new recrystallized grains are produced. As shown in Fig. 2, the width of the complex stacking faults is very large because of the low stacking fault energy in the sample of our experiment. This situation is disadvantageous for D R C which requires dislocation climbing, and these conditions suppress D R C and
induce DRX. It should be noted that the formation of networks and the misorientation development between those cells play important roles in DRX. In the sample compressed at 1650 K, rectangular networks cannot be seen clearly. The width of the complex stacking faults is narrower than that at 1490 K and subboundaries were observed as shown in Figs. 6 and 7. This means that the stacking fault energy
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phase in a two-phase stainless steel refined the grain size of the a phase and then the resultant grain boundary sliding caused superplasticity. On the other hand, Takasugi et al. [18, 19] indicated that grain boundary sliding is a main deformation mechanism in superplastic Ni3(Si,Ti ). In Zr3AI [4] and (Fe, Co)3V [12], the grain boundary sliding of DRX grains was found to lead to ductilization. In these experiments, grain boundary sliding was confirmed by the large misorientation between adjacent grains and the existence of dislocation-free grains. The mixed structure of D R X grains with and without dislocations can be seen in the sample compressed at 1490 K to the strain of 0.33 in our experiment. In addition, there exist large grains with a high dislocation density which are considered to be initial grains. This indicates that none of the initial grains has been replaced by D R X grains, although D R X grains are predominant. Thus, it is concluded that the ductile deformation in Nb3AI polycrystal is due to the occurrence of DRX at initial grain boundaries and subsequent grain boundary sliding.
5. Summarizing remarks
Fig. 8. Dislocations near a grain boundary of a sample compressed at 1490 K to a strain of 0.07. increases with increasing temperature. New grains formed by DRX were observed less frequently at grain boundaries in the sample compressed at 1650 K than in that compressed at 1490 K. These microstructural observa.tions are consistent with the fact that, as the compression temperature increases, the amount of deformation softening decreases. The sample compressed at 1650 and 1800 K did not exhibit a clear peak stress. The flow stress shifts immediately from the stage of initial strain hardening to steady state deformation. This is because the restoration process occurs rapidly at these temperatures. 4.2. Role o f D R X f o r ductifity
Since superplasticity was found in Fe3(AI , Si) [5, 6], superplasticity has been reported in many intermetallic compounds [10, 15, 16, 18-20, 26-28]. Superplasticity in most intermetallic compounds has been discussed in terms of DRX. Recently, D R X has been found to play an important role in the superplasticity of some disordered alloys. Maehara [29] proposed that superplasticity is caused by the balance between D R X and the strain hardening in a duplex phase stainless steel. Tsuzaki et al. [30] suggested that D R X of the a
Off-stoichiometric, single-phase Nb3AI produced by powder metallurgy deformed plastically above 1400 K and a ductile deformation behaviour without macroscopic cracking appeared above 1490 K. The stress-strain curves of the sample showed deformation softening owing to DRX. It is revealed that the formation of complex stacking faults bounded by partial dislocations and characteristic dislocation networks and subboundaries results in DRX.
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