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Effect of heterogeneous precipitation on age-hardening of Al2O3 particle dispersion Al-4mass%Cu composite produced by mechanical alloying
Scripta mater. 42 (2000) 755–760 www.elsevier.com/locate/scriptamat
EFFECT OF HETEROGENEOUS PRECIPITATION ON AGEHARDENING OF Al2O3 PARTICLE DISPERSION Al-4mass%Cu COMPOSITE PRODUCED BY MECHANICAL ALLOYING S. Arakawa, T. Hatayama, K. Matsugi and O. Yanagisawa Faculty of Engineering, Hiroshima University, 1-4-1 Kagamiyama, Higashi-Hiroshima, Hiroshima, 739-8527, Japan (Received December 1, 1999) (Accepted December 14, 1999) Keywords: Mechanical alloying; Aluminum-copper alloy; Composite; Aging
Introduction The acceleration of aging kinetics has been frequently observed in aluminum matrix composites produced by ingot or powder metallurgy (1–3). The most probable reason is ascribed to high dislocation density in the vicinity of ceramic particles due to the difference in thermal expansion coefficient between the matrix and ceramic (4,5). The higher dislocation density enhances the diffusion of solute atoms, and serves greater preferential precipitation sites, which leads to hasten the nucleation and growth of the precipitates, and consequently, the composites reach the peak hardness in shorter time than the materials without reinforcements (2,6). Recently, in the mechanically alloyed (MA) Al-4mass%Cu/Al2O3 composites, the authors have found that the age-hardening response significantly decreases, and that considerable stable phases are formed at a very short aging time (7). MA materials have not only the high dislocation density around ceramic particles but also ultrafine grain structures of submicron size (8,9). Grain boundaries are also preferential precipitation sites that increase with decreasing in a grain size. The age-hardenability will decrease if the precipitation of stable phases on grain boundaries increases in quantity, because the stable phases contribute to hardening very little, and reduce solute atoms in the matrix. The purposes of this study are to investigate the local precipitation behaviors, and attempt to clarify the dominant microstructural factors of the decrease in the age-hardenability and the acceleration of the age-hardening kinetics in the Al2O3 particle dispersion Al-4mass%Cu composites produced by mechanical alloying. In order to build a basis for comparison, the age-hardening behaviors of the unreinforced matrix alloy (IM alloy), which is produced by ingot metallurgy technique, are also investigated.
Experimental Procedures The elemental powders were pure aluminum (purity: 99.9 mass%, average size: 22 m) and copper (99.99 mass%, 10 m). These powders were blended to the Al-4mass%Cu composition in the matrix. The mixed powders were prepared by mechanical alloying using a planetary ball mill in argon atmosphere. The stearic acid of 1 mass% was charged to the powders as the process-control agent. The 1359-6462/00/$–see front matter. © 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(99)00426-1
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Figure 1. Aging curves of the Al-4mass%Cu/10.4vol%Al2O3 composite (MA: open marks) and the Al-4mass%Cu alloy (IM: closed marks) aged at various temperatures.
contamination from a mill pot and balls through wear have been observed in milled powders (7), so an alumina (␣-Al2O3) mill pot and balls were used here. The alumina pot and balls are very brittle, so the irregular Al2O3 particles by the wear mingled with the milled powders in the milling processes. In this study, the dispersed Al2O3 particles in the matrix were supplied from the mill pot and balls, and the Al2O3 particle content was controlled by the milling time. The MA powders milled for 64.8 ks were cold pressed to form compacts. The compacts were hot-extruded (extrusion ratio ⫽ 9) at 823K after degassing at 673K for 7.2 ks in vacuum, when the average size and volume fraction of Al2O3 particles were 2.7 m and 10.4 vol%, respectively. The extruded bars were machined to the specimens with the length of 10 mm and the diameter of 6 mm. The specimens were held at 813K for 7.2 ks, and then quenched into iced water. Aging in an oil bath was carried out immediately at 393– 463K after quenching. The average grain sizes of MA and IM materials after quenching were 0.34 and 597 m, respectively. The hardness were measured by Vickers microhardness test. X-ray diffraction analysis, scanning (SEM) and transmission electron microscope (TEM) observations were performed to investigate the microstructures. Results The age-hardening curves of the MA composite and the IM alloy at various temperatures are shown in Fig. 1. It is noted that the MA composite aged at 393K shows clear age-hardening and the maximum peak of hardness appears at extremely shorter time (86.4 ks) compared with the IM alloy (about 6000 ks). The age-hardenability of the MA composite, ⌬HV (the amount of age-hardening), is, however, only the ⌬HV of 26, which is about a half of the value of the IM alloy (⌬HV ⫽ 50). The MA composites show no age-hardening above 433K where the hardness decreases with aging time (age-softening), nevertheless the IM alloy at 463K has a clear peak. Figure 2(a) and (b) show a higher density of dislocation in the vicinity of Al2O3 particle after quenching and the precipitates on grain boundaries after aging at 463K for 1.8 ks, respectively. There are many phases on grain boundaries at the early stage of aging and little in grains as shown in Fig. 2(b). The precipitation of phases was dependent on a location and some grain boundaries contained no precipitates. The reflection electron images of the microstructures in the MA composites aged at 393 and 463K are shown in Fig. 3. White and large gray particles in the microstructures are corresponded to the phases and Al2O3 particles, respectively. The phases become more clear and large with increasing aging time. It seems from the figure that the phases are classified to the two groups: (1) the phases marked by A in Fig. 3 (A) that precipitate in the network-like morphology; (2) the phases
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Figure 2. TEM micrographs of the Al-4mass%Cu/10.4vol%Al2O3 MA composite after (a) quenching and (b) aging at 463K for 1.8 ks.
marked by B in Fig. 3 (B) that are smaller than A. The A phases appeared on grain boundaries at first, and then the B phases mainly within grains in process of aging time. Figure 2 and 3 indicate that the fair amount of phases preferentially precipitates on grain boundaries. It is of interest that the phases are generally recognized at a short distance from Al2O3 particles, and little at the matrix/Al2O3 interfaces. Figure 4 shows the number density of the phases and the ratio of integrated intensity of phase to the matrix. It is seen that the phase increases with aging time, and that comparing Fig. 1 and 4, the considerable quantity of phases precipitates at the peak hardness. The change of number density can be easily divided into the five regions; I, II, III, IV and V. In the region I, the A phases preferentially precipitated on grain boundaries and increased the number density with aging time. In region II, the growth of A phases occurred and the B phases began to precipitate, however, the growth of A phases was dominant, so that the number density decreased. In the region III and IV, the number density of the phases increased, because the precipitation rate of the B phases became dominant compared with the
Figure 3. Reflection electron images of the phases in the Al-4mass%Cu/10.4vol%Al2O3 MA composite aged at 393K ((a) after aging for 86.4 ks, (b) 532.8 ks) and 463K ((c) 3.6 ks, (d) 172.8 ks).
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Figure 4. Number densities and ratios of integrated intensity of the phase to the matrix in the Al-4mass%Cu/10.4vol%Al2O3 MA composites aged at 393 and 463K as a function of aging time.
growth rate of the A phases. The growth of the A and B phases was dominant in the region V, and the number density decreased again. Figure 5 shows the average size and the size distribution of phases during aging. The size and the standard deviation of phases increase with aging time. The average size at 463K is larger than that at 393K in any aging time. The size distribution shows a single sharp peak in the region I and II as shown in Fig. 5(a) and (c), and separates to the two clear peaks in region IV as shown in Fig. 5(b) and (d). The peaks, PA and PB, correspond to those of the A and B phases, respectively. Discussion It is well known that the precipitation processes of Al-Cu alloys complete as following sequence; GP zone 3 GPII zone 3 ⬘ intermediate 3 (CuAl2) phase (10). The GPII zone and ⬘ intermediates are most effective for the increase of hardness, and the precipitation of the stable phase decreases hardness (over aging). The MA composite aged at 393K shows clear age-hardening peak at 86.4 ks (Fig. 1), where the precipitation of phases occurs at the same time (Fig. 4 and 5). Grain boundaries are paths of high diffusivity and preferential precipitation sites, where the precipitation rate and the growth rate are very high compared with the rate within a grain, and stable phases appear in the early stage of aging process (11,12). In the common materials with large grains,
Figure 5. Size distribution of phases in the Al-4mass%Cu/10.4vol%Al2O3 MA composites aged at 393 and 463K, where d and S.D. are the average size and standard deviation of phase, respectively.
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the precipitates on grain boundaries are very small, so that the precipitates could not give a significant effect to the age-hardening behavior. In contrast, the MA composites have an ultrafine grain structure of order of submicron, the area of the grain boundaries is large, and the volume of the precipitates, A, not contributing to age-hardening, extends to fair amount, which significantly decreases the agehardenability due to decreasing solutes in the matrix. In the area of a higher dislocation density, the ⬘ intermediate can be considered to precipitate at first, and then to change into phase. The simultaneous change and the growth of the phases in the area of a higher dislocation density from the early stage of aging also decrease age-hardenability, because the peak hardness becomes high when the solute atoms precipitate out in the matrix on a fine scale. In general, the time to the peak hardness extends as the concentration of solute atoms in the matrix decreases (10). The precipitation of phases at the early stage of aging decreases the concentration of Cu in the matrix, which predicts to increase the time to the peak hardness. The peak time of the MA composites, however, remarkably decreases compared with that of the IM alloy (Fig. 1), and the peak hardness aged at 393K appears at the short time of 86.4 ks, nevertheless the fair amount of phase precipitates at the same time (Fig. 4). The apparent activation energy of the aging processes responsible for the peak hardness was 51 kJ/mol in MA composites, and 118 kJ/mol in IM alloys. The earlier attainment of the peak hardness can be ascribed to the precipitation assisted by dislocations, because the diffusivity of solute atoms would be dominated by dislocations and grain boundary diffusion in this temperature, and the precipitates on grain boundaries can be considered to contribute to the peak hardness very little (13). In the temperature above 433K, the change of the intermediates on the dislocation into phases and the growth of phases become more active (Fig. 4), and the MA composites could not be strengthened by the precipitates, so that age-softening (over aging state) occurs from the initial stage of aging at 463K as shown in Fig. 1. The grain size is very small in the MA composite and consequently diffusion distance to grain boundaries is very short, so the solute atoms can easily reach grain boundaries with increasing temperature, which would also increase the quantity of phases on grain boundaries (Fig. 2 and 4).
Conclusion Mechanically alloyed Al-4mass%Cu composites with Al2O3 particle 10.4 vol% have very fine grains of 0.34 m and a higher dislocation density in the vicinity of the Al2O3 particles. In the MA composite, the phases are generally recognized at a short distance from Al2O3 particle, and little at the matrix/Al2O3 interfaces. The age-hardenability of the MA composite significantly decreases compared with the IM alloy, because the preferential precipitation of the stable phases on grain boundaries occurs in an earlier stage of aging, and extends to fair amount due to very fine grains, which reduces the concentration of copper atoms in the matrix to form GP zone and intermediates. There is remarkable acceleration of aging kinetics in the MA composite, which is ascribed to the precipitation of intermediates assisted by the higher dislocation density.
References 1. 2. 3. 4. 5. 6.
T. G. Nieh and R. F. Karlak, Scripta Metall. 18, 25 (1984). Y. Song and T. N. Baker, Mater. Sci. Eng. 10, 406 (1994). G. M. Janowski and B. J. Pletka, Metall. Trans. 26A, 25 (1995). M. Vogelsang, R. J. Arsenault, and R. M. Fisher, Metall. Trans. 17A, 379 (1986). I. Dutta and D. L. Bourell, Mater. Sci. Eng. A112, 813 (1989). T. Christman and S. Suresh, Acta Metall. 36, 1691 (1988).
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7. 8. 9. 10. 11. 12. 13.
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S. Arakawa, T. Hatayama, K. Matsugi, and O. Yanagisawa, Proceedings of the 3rd Asia-Pacific Symposium on Advances in Engineering Plasticity and Its Applications, p. 457 (1996). J. S. Benjamin and M. J. Bomford, Metall. Trans. 8A, 1301 (1977). P. S. Gilma and W. D. Nix, Metall. Trans. 12A, 813 (1981). J. M. Silcock, T. J. Heal, and H. K. Hardy, J. Inst. Metals. 82, 239 (1953–54). M. B. Chamberlain and S. L. Lehocsky, Thin Solid Films 45, 189 (1977). D. R. Frear, J. E. Sanchez, A. D. Romig, Jr., and J. W. Morris, Jr., Metall. Trans. 21A, 2449 (1990). S. K. Hong, J. C. Choi, D. Z. Lin, H. Tezuka, T. Sato, and A. Kamio, J. Jpn. Inst. Metals. 60, 569 (1996).
EFFECT OF HETEROGENEOUS PRECIPITATION ON AGEHARDENING OF Al2O3 PARTICLE DISPERSION Al-4mass%Cu COMPOSITE PRODUCED BY MECHANICAL ALLOYING S. Arakawa, T. Hatayama, K. Matsugi and O. Yanagisawa Faculty of Engineering, Hiroshima University, 1-4-1 Kagamiyama, Higashi-Hiroshima, Hiroshima, 739-8527, Japan (Received December 1, 1999) (Accepted December 14, 1999) Keywords: Mechanical alloying; Aluminum-copper alloy; Composite; Aging
Introduction The acceleration of aging kinetics has been frequently observed in aluminum matrix composites produced by ingot or powder metallurgy (1–3). The most probable reason is ascribed to high dislocation density in the vicinity of ceramic particles due to the difference in thermal expansion coefficient between the matrix and ceramic (4,5). The higher dislocation density enhances the diffusion of solute atoms, and serves greater preferential precipitation sites, which leads to hasten the nucleation and growth of the precipitates, and consequently, the composites reach the peak hardness in shorter time than the materials without reinforcements (2,6). Recently, in the mechanically alloyed (MA) Al-4mass%Cu/Al2O3 composites, the authors have found that the age-hardening response significantly decreases, and that considerable stable phases are formed at a very short aging time (7). MA materials have not only the high dislocation density around ceramic particles but also ultrafine grain structures of submicron size (8,9). Grain boundaries are also preferential precipitation sites that increase with decreasing in a grain size. The age-hardenability will decrease if the precipitation of stable phases on grain boundaries increases in quantity, because the stable phases contribute to hardening very little, and reduce solute atoms in the matrix. The purposes of this study are to investigate the local precipitation behaviors, and attempt to clarify the dominant microstructural factors of the decrease in the age-hardenability and the acceleration of the age-hardening kinetics in the Al2O3 particle dispersion Al-4mass%Cu composites produced by mechanical alloying. In order to build a basis for comparison, the age-hardening behaviors of the unreinforced matrix alloy (IM alloy), which is produced by ingot metallurgy technique, are also investigated.
Experimental Procedures The elemental powders were pure aluminum (purity: 99.9 mass%, average size: 22 m) and copper (99.99 mass%, 10 m). These powders were blended to the Al-4mass%Cu composition in the matrix. The mixed powders were prepared by mechanical alloying using a planetary ball mill in argon atmosphere. The stearic acid of 1 mass% was charged to the powders as the process-control agent. The 1359-6462/00/$–see front matter. © 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S1359-6462(99)00426-1
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Figure 1. Aging curves of the Al-4mass%Cu/10.4vol%Al2O3 composite (MA: open marks) and the Al-4mass%Cu alloy (IM: closed marks) aged at various temperatures.
contamination from a mill pot and balls through wear have been observed in milled powders (7), so an alumina (␣-Al2O3) mill pot and balls were used here. The alumina pot and balls are very brittle, so the irregular Al2O3 particles by the wear mingled with the milled powders in the milling processes. In this study, the dispersed Al2O3 particles in the matrix were supplied from the mill pot and balls, and the Al2O3 particle content was controlled by the milling time. The MA powders milled for 64.8 ks were cold pressed to form compacts. The compacts were hot-extruded (extrusion ratio ⫽ 9) at 823K after degassing at 673K for 7.2 ks in vacuum, when the average size and volume fraction of Al2O3 particles were 2.7 m and 10.4 vol%, respectively. The extruded bars were machined to the specimens with the length of 10 mm and the diameter of 6 mm. The specimens were held at 813K for 7.2 ks, and then quenched into iced water. Aging in an oil bath was carried out immediately at 393– 463K after quenching. The average grain sizes of MA and IM materials after quenching were 0.34 and 597 m, respectively. The hardness were measured by Vickers microhardness test. X-ray diffraction analysis, scanning (SEM) and transmission electron microscope (TEM) observations were performed to investigate the microstructures. Results The age-hardening curves of the MA composite and the IM alloy at various temperatures are shown in Fig. 1. It is noted that the MA composite aged at 393K shows clear age-hardening and the maximum peak of hardness appears at extremely shorter time (86.4 ks) compared with the IM alloy (about 6000 ks). The age-hardenability of the MA composite, ⌬HV (the amount of age-hardening), is, however, only the ⌬HV of 26, which is about a half of the value of the IM alloy (⌬HV ⫽ 50). The MA composites show no age-hardening above 433K where the hardness decreases with aging time (age-softening), nevertheless the IM alloy at 463K has a clear peak. Figure 2(a) and (b) show a higher density of dislocation in the vicinity of Al2O3 particle after quenching and the precipitates on grain boundaries after aging at 463K for 1.8 ks, respectively. There are many phases on grain boundaries at the early stage of aging and little in grains as shown in Fig. 2(b). The precipitation of phases was dependent on a location and some grain boundaries contained no precipitates. The reflection electron images of the microstructures in the MA composites aged at 393 and 463K are shown in Fig. 3. White and large gray particles in the microstructures are corresponded to the phases and Al2O3 particles, respectively. The phases become more clear and large with increasing aging time. It seems from the figure that the phases are classified to the two groups: (1) the phases marked by A in Fig. 3 (A) that precipitate in the network-like morphology; (2) the phases
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Figure 2. TEM micrographs of the Al-4mass%Cu/10.4vol%Al2O3 MA composite after (a) quenching and (b) aging at 463K for 1.8 ks.
marked by B in Fig. 3 (B) that are smaller than A. The A phases appeared on grain boundaries at first, and then the B phases mainly within grains in process of aging time. Figure 2 and 3 indicate that the fair amount of phases preferentially precipitates on grain boundaries. It is of interest that the phases are generally recognized at a short distance from Al2O3 particles, and little at the matrix/Al2O3 interfaces. Figure 4 shows the number density of the phases and the ratio of integrated intensity of phase to the matrix. It is seen that the phase increases with aging time, and that comparing Fig. 1 and 4, the considerable quantity of phases precipitates at the peak hardness. The change of number density can be easily divided into the five regions; I, II, III, IV and V. In the region I, the A phases preferentially precipitated on grain boundaries and increased the number density with aging time. In region II, the growth of A phases occurred and the B phases began to precipitate, however, the growth of A phases was dominant, so that the number density decreased. In the region III and IV, the number density of the phases increased, because the precipitation rate of the B phases became dominant compared with the
Figure 3. Reflection electron images of the phases in the Al-4mass%Cu/10.4vol%Al2O3 MA composite aged at 393K ((a) after aging for 86.4 ks, (b) 532.8 ks) and 463K ((c) 3.6 ks, (d) 172.8 ks).
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Figure 4. Number densities and ratios of integrated intensity of the phase to the matrix in the Al-4mass%Cu/10.4vol%Al2O3 MA composites aged at 393 and 463K as a function of aging time.
growth rate of the A phases. The growth of the A and B phases was dominant in the region V, and the number density decreased again. Figure 5 shows the average size and the size distribution of phases during aging. The size and the standard deviation of phases increase with aging time. The average size at 463K is larger than that at 393K in any aging time. The size distribution shows a single sharp peak in the region I and II as shown in Fig. 5(a) and (c), and separates to the two clear peaks in region IV as shown in Fig. 5(b) and (d). The peaks, PA and PB, correspond to those of the A and B phases, respectively. Discussion It is well known that the precipitation processes of Al-Cu alloys complete as following sequence; GP zone 3 GPII zone 3 ⬘ intermediate 3 (CuAl2) phase (10). The GPII zone and ⬘ intermediates are most effective for the increase of hardness, and the precipitation of the stable phase decreases hardness (over aging). The MA composite aged at 393K shows clear age-hardening peak at 86.4 ks (Fig. 1), where the precipitation of phases occurs at the same time (Fig. 4 and 5). Grain boundaries are paths of high diffusivity and preferential precipitation sites, where the precipitation rate and the growth rate are very high compared with the rate within a grain, and stable phases appear in the early stage of aging process (11,12). In the common materials with large grains,
Figure 5. Size distribution of phases in the Al-4mass%Cu/10.4vol%Al2O3 MA composites aged at 393 and 463K, where d and S.D. are the average size and standard deviation of phase, respectively.
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the precipitates on grain boundaries are very small, so that the precipitates could not give a significant effect to the age-hardening behavior. In contrast, the MA composites have an ultrafine grain structure of order of submicron, the area of the grain boundaries is large, and the volume of the precipitates, A, not contributing to age-hardening, extends to fair amount, which significantly decreases the agehardenability due to decreasing solutes in the matrix. In the area of a higher dislocation density, the ⬘ intermediate can be considered to precipitate at first, and then to change into phase. The simultaneous change and the growth of the phases in the area of a higher dislocation density from the early stage of aging also decrease age-hardenability, because the peak hardness becomes high when the solute atoms precipitate out in the matrix on a fine scale. In general, the time to the peak hardness extends as the concentration of solute atoms in the matrix decreases (10). The precipitation of phases at the early stage of aging decreases the concentration of Cu in the matrix, which predicts to increase the time to the peak hardness. The peak time of the MA composites, however, remarkably decreases compared with that of the IM alloy (Fig. 1), and the peak hardness aged at 393K appears at the short time of 86.4 ks, nevertheless the fair amount of phase precipitates at the same time (Fig. 4). The apparent activation energy of the aging processes responsible for the peak hardness was 51 kJ/mol in MA composites, and 118 kJ/mol in IM alloys. The earlier attainment of the peak hardness can be ascribed to the precipitation assisted by dislocations, because the diffusivity of solute atoms would be dominated by dislocations and grain boundary diffusion in this temperature, and the precipitates on grain boundaries can be considered to contribute to the peak hardness very little (13). In the temperature above 433K, the change of the intermediates on the dislocation into phases and the growth of phases become more active (Fig. 4), and the MA composites could not be strengthened by the precipitates, so that age-softening (over aging state) occurs from the initial stage of aging at 463K as shown in Fig. 1. The grain size is very small in the MA composite and consequently diffusion distance to grain boundaries is very short, so the solute atoms can easily reach grain boundaries with increasing temperature, which would also increase the quantity of phases on grain boundaries (Fig. 2 and 4).
Conclusion Mechanically alloyed Al-4mass%Cu composites with Al2O3 particle 10.4 vol% have very fine grains of 0.34 m and a higher dislocation density in the vicinity of the Al2O3 particles. In the MA composite, the phases are generally recognized at a short distance from Al2O3 particle, and little at the matrix/Al2O3 interfaces. The age-hardenability of the MA composite significantly decreases compared with the IM alloy, because the preferential precipitation of the stable phases on grain boundaries occurs in an earlier stage of aging, and extends to fair amount due to very fine grains, which reduces the concentration of copper atoms in the matrix to form GP zone and intermediates. There is remarkable acceleration of aging kinetics in the MA composite, which is ascribed to the precipitation of intermediates assisted by the higher dislocation density.
References 1. 2. 3. 4. 5. 6.
T. G. Nieh and R. F. Karlak, Scripta Metall. 18, 25 (1984). Y. Song and T. N. Baker, Mater. Sci. Eng. 10, 406 (1994). G. M. Janowski and B. J. Pletka, Metall. Trans. 26A, 25 (1995). M. Vogelsang, R. J. Arsenault, and R. M. Fisher, Metall. Trans. 17A, 379 (1986). I. Dutta and D. L. Bourell, Mater. Sci. Eng. A112, 813 (1989). T. Christman and S. Suresh, Acta Metall. 36, 1691 (1988).
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7. 8. 9. 10. 11. 12. 13.
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S. Arakawa, T. Hatayama, K. Matsugi, and O. Yanagisawa, Proceedings of the 3rd Asia-Pacific Symposium on Advances in Engineering Plasticity and Its Applications, p. 457 (1996). J. S. Benjamin and M. J. Bomford, Metall. Trans. 8A, 1301 (1977). P. S. Gilma and W. D. Nix, Metall. Trans. 12A, 813 (1981). J. M. Silcock, T. J. Heal, and H. K. Hardy, J. Inst. Metals. 82, 239 (1953–54). M. B. Chamberlain and S. L. Lehocsky, Thin Solid Films 45, 189 (1977). D. R. Frear, J. E. Sanchez, A. D. Romig, Jr., and J. W. Morris, Jr., Metall. Trans. 21A, 2449 (1990). S. K. Hong, J. C. Choi, D. Z. Lin, H. Tezuka, T. Sato, and A. Kamio, J. Jpn. Inst. Metals. 60, 569 (1996).