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Sintering of nanostructured W-Cu alloys prepared by mechanical alloying
NmoStru~
Peqpmon
Matmiek.Vol. 10. No. 2, pi. 283-290.1998 Ekevia ScienceLtd 8 1998AC&Mctdlurgi~ Inc. FvinkdintheUSA. All ri8ht.s ralavd 096%9773/98 $19.00 + .OO
PKISO9659773(98)00065-S
SINTERING OF NANOSTRUCTURED W-Cu ALLOYS PREPARED BY MECHANICAL ALLOYING Jin-Chun Kim and In-Hyung Moon Department of Materials Engineering, Hanyang University, Seoul 133-791, Korea (Accepted October 2,1997)
Absbwt - Nanwtructured (NS) powders with compositions corresponding to W2Owt%Cu and W-3Owt%Cu were prepared by mechanical alloying. The microstructure and grain size of as-milled and annealedpowders were analyzed by transmission electron microscopy. The compacted specimens were sintered at temperatures in the range lOOO”-13&I”, and then the microstructures of sintered parts were analyzed by scanning electron microscopy. Sintering of mechanicaily alloyed W-Cu alloysappears to be independent of Cu content, and may be explained in terms of recovery and grain growth in the mechanically alloyed powders as well as impurity activated simering of W. After sintering. Cu pools areformed outside the mechanically alloyed powders. A relative sintered density of more than 95% is obtained by particle rearrangement during liquid!-phase sintering, and the greatest homogeneity of W and Cu phases is achieved by sintering at i’2009 01998 Acta Metallurgica Inc.
INTRODUCTION Recendy, W-Cu alloys have been used for thermal managing and as microwave materials due to the high thermal conductivity of copper and the low thermal expansion coefficient of tungsten (l-3). The enhanced thermal properties of these alloys can be achieved by the unique combination of elemental phases. Research has been conducted on improving the homogeneity of component phases in these alloys. Among them, the use of nanostructured W-Cu powders prepared by coreduction of oxide powders (45) and mechanical alloying (MA) of elemental powders (6,;‘) has been widely studied. We investigated the characteristics of nanostructured W-Cu alloys prepared by MA (8). In this experiment, it was found that the grain growth of the MA powders significantly occurred at solid-state sintering temperatures below 1083O. The internal structure of MA powders after sintering was similar to the conventional one of W-Cu alloys processed by liquid-phase s&ring. The enhanced sinterability of MA W-Cu alloys was also obtained during solid-state sintering. After sintcring, a high sintered density and a homogeneity of W-Cu alloys were obtained. In the present work, the sintering behavior of MA W-Cu alloys is described and a possible explanation for the observations is suggested. 283
J-C KIM AND1-H kGN
EXPERIMENTAL
Elemental W (99.9% purity) and Cu (99.5% purity) powders with particle size of 4.28 l.trn and 50.42 p.m, respectively, were used in this experiment. Composite powders with compositions corresponding to the W-20wt%Cu and W-30wt%Cu were prepared by mechanical alloying. The MA was performed in an attritor of 750 cm3 capacity using stainless steel balls with 0.48 cm diameter. The weight ratio of ball to powder was maintained to60: 1. The microstructure and grain size of as-milled powders and annealed powders at 900” were analyzed by transmission electron microscopy (TEM, model:2000-EX II, JEOL) by attaching the powders on the 300 mesh grid. The bulk specimens of MA powders were prepared by cold pressing into a disk shape with a green density of 40&2% of theoretical density. The compacted MAW-3OCu specimens were sintered in the temperature range looO”-1300” for 1 hr in HZatm. The sintering of the MAW-2OCu compacts was carried out at 1000° and 1050’ for 0 minute in order to investigate the microstructural change of specimens during the heat-up stage( 0 min. means that the holding time at the sintering temperature is zero). The microstructures of sintered parts were analyzed by scanning electron microscopy (SEM).
RESULTS
In mechanical alloying, the development of particle shape and size reaches steady-state through several different stages of processing, as suggested by Benjamin (9). From the observation of morphological development and change of x-ray diffraction patterns with MAtimes, the milling time of the steady-state condition in this experiment was 100 hrs. In that case, the MA powders were characterized by three dimensional equiaxed shape, and its size was 3-5 l.trn.
Figure 1. The TEM micrographs of the mechanically alloyed W-30wt%Cu powder for 100 hr; bright field TEM image (a) and selected area diffraction (SAD) pattern (b).
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Figure 1 shows the bright field TEM image and selected area diffraction (SAD) pattern of MA W-30& powders at the steady-state condition. As shown in this figure, the grain size of MA W-Cu powders is 20 to 30 nm, and the SAD pattern of matrix with sharp diffracted rings indicates that the W and Cu remain in the crystalline state. In addition, the irregular grain morphology of the as-milled powder was observed and W phases are not distinguishable from the Cu phases. Figure 2 shows the SEM and TEM microstructures of the MA W-3001 powders annealed at 900” in Hz atmosphere. As shown in these figures, the grain size of MApowders (in Figure 2(b)) with annealing for 5 hrs increases from 20-30 nm of as-milled powders to about 130 nm. Now, we are able to distinguish W from Cu in the annealed MA powders, and it shows that the W grains undergo shape change from irregular shape to spheroid. In addition, some of the W-grains are coalesced or even sintered (marked with a circle). They exhibited similar microstructural properties fa.bricated by the conventional liquid-phase sintering in the W-Cu system using microsized powers (10.11). This coalescence of W/W grains is clearly visible in the TEM microstructure (Figure 2(c):). In this figure, a neck formation due to W-W sintering is evident (arrow A), and also W grain is bounded by the Cu phase (arrow B) and surrounded by the Cu layer. However, the external morphology of MA powder (Figure 2(a)) is almost identical to that of initial MA powder regardless of different thermal treatment. Figure 3 shows the SEM microstructures of MA W-3OCu compact sintered at 1000” for 1 hr. The W-grain size of the MA powders increases to about 200 nm at this temperature. The interior microstructure of the MA W-Cu powders observed with SEM exhibits a typical microstructure of fully densified W-Cu alloys by liquid phase sintering or Cu infiltration method (12). It is also observed that the Cu rich phase or pools indicated by the arrow covers the exterior of MA powders. The morphology of MA powders in the sintered part (Figure 3(a)) is similar to that of initial MA
Figure 2. The SEM micrographs of powder (a) and internal structure (b) and bright field TEM micrograph (c) of the MA W-30wt%Cu powders annealed at 900” in H2 atm. for 5 hr.
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powder and annealed powder at 900°, and the MA powders are aggregated by the Cu pools. The details on the formation of Cu pools at solid-stage sintering will be discussed in the separate paper (13). As shown in Figures 2 and 3, the grain size of the MA powders drastically increases from 20-30 nm to 200 nm at the solid-stage sintering stage and theinternal microstructure of MA powder which was formed from the coalescence or sintering of W grains is similar to that of typical liquidphase sintering of the W-Cu alloys. If the sintering of the MA W-Cu powders takes place, the Wgrain size of the MA powders drastically increases from the initially fine-grained as-milled powders withoutchanging thepowdershapeandtheintemai microstmctureofMApowderis very similar to that of the liquid-phase sintering, and notably the Cu pools are formed outside MA
Figure 3. The SEM micrographs of the MA W-3Owt%Cu sintered at 1000’ for 1 hr in Hz atm.; Cu pool area (a) and interior area of powder (b).
Figure 4. The SEM micrographs of the MA W-20wtlCu sintered for 0 min in HZ atm. at various temperatures; 1000° (a) and 1050’ (b).
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Figure ,4(a) and (b) exhibit the microstructure of MA W-20wt%Cu specimens sintered at 1000” and 10!jO”for0 min in HZ atm,respectively. As shown in these figures, sintering also occurs, and its shucture is well matched to that of the MA W-3OCu specimen. Based on this observation, we suggest thlat the nanosintering of MA W-Cu powders may be independent of Cu content, and we find its origin in the thermal effects of the heating-up stage. Figure 5 shows SEM micrographs of MA W-3Owt%Cu specimens sintered at 1100” (a), 1150” (b), 12OOO(c)and13OO”(d)for 1 hrinH2atm. TheW-grainsizeof8OOto lOOOnm,asshown in these figums, is much larger compared to that of the specimens sintered at solid-state sintering, and the mixing state of W and Cu phase is more homogeneous with increasing sintering temperature. The sintered density of all specimens is higher than the relative density of 95%, and there is no observable porosity. The MA powder shape in the specimen sintered at 1100’ is not changed fmm the initial morphology of MA powders, but the W grains are totally surrounded by the Cu liquid phase. At 1150’ (in Figure 5(b)), the shape of agglomerated W grains begins to change to the homogeneous state without further growth of the W grains. In the case of the specimen sir&red at 1200’ (Figure 5(c)), it shows that each W grain in the MA powder is rearranged, alnd the highest homogenous state of W and Cu phases can be attained by sintering. Finally, the microstructure of the specimen sintered at 1300” is observed to be analogous to that of the liquid-phase sintering investigated by others, as shown in Figure 5(d).
DISCUSSION 1. Solid-stateSintering In general, the grain growth of the W-Cu alloys during the solid-state sintering is negligible because W and Cu are completely immiscible. However, as described above, the grain size of MA powders drasticallyincreases during the solid-state sintering and its inner microstructure gradually changes to tlhat of conventional liquid-phase sintering. This phenomenon may be explained in terms of following mechanisms. Firstly, it may be derived from the recovery and grain growth due to internal strain energy, which is caused by the repeated cold welding and fracture during mechanical alloying process. The stored energy of the MA powders will be released in the form of external thermal effect during heating (14,15). Consequently, it can be assumed that the unstable W and Cu grains and lattice deformations in the MA powders transform into the stable state with increasing temperature, i.e., the recovery and recrystallization occur during the heating, and then the grain growth occurs in the MA powders. The second possibility is the activated sintering of W due to the Fe impurity, which is inherent in the MA powders from the ball media and attritor impeller during the MA process. The effect of impurity addition such as Fe, Ni and Co etc., which causes the activated sintering of W, was well known to enhance the sintering behavior of W (16,17). Third,. the activated liquid phase sintering of W-Cu due to impurity must be considered (10,18,19). The addition of the activators which are soluble in the liquid phase enhances densification and grain growth through solution and reprecipitation. However, it should be carefully colnsidered that the impurity effect on W-Cu alloys only acts at the liquid-phase sintering and it shouldbe accompanied with the continuous grain growth in the liquid phase sintering stage.
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Figure 5. The micrographs observed by the MA W-30wt%Cu sintered for 1 hr at various temperatures: 1100’ (a), 1150 (b), 1200 (c) and 1300’ (d). In our study, the growth of W grain is solely observed during the solid-state sintering, and no significant grain growth occurs in the liquid-phase sintering. Finally, we should consider the interreaction of the Wand the Cu constituents. Under the equilibrium condition, W and Cu are completely immiscible due to the high positive heat of mixing; 35SKJ/mol(20). But, Panichkina (21) reported that the considerable grain growth of W particle was observed with increasing W solubility in Cu, which resulted from the use of W particle of sub-micron size. However, this mechanism cannot be directly applied to the W-grain growth in the MA powders at the solid-state sintering because his experiment was carried out only in the range of the liquid-phase sintering. Besides the above mechanisms, the solubility increase of Cu in W due to the MA as well as the decrease of Cu melting temperature due to the Gibbs-Thomson effect should be taken into account. Among them, the mechanism of recovery followed by grain growth of the MA powder and the impurity activated sintering of W seem to be the case in this experiment. No matter what mechanisms operate, however, the drastic grain growth in the NS W-Cu powders prepared by mechanical alloying, i.e., nanosintering occurred, and the Cu pools were formed in the external of the MA powders.
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2. Microstructure Changes of Liquid Phase Sintering Stage
In this experiment, the relative sintered density of the nanostructured W-Cu alloys increased drastically above 95% in the solid-stage sintering stage at 1050’ and the high sintered density of 98% wasobtainedat liquid-phase sintering temperatures (at 1100°). This result maybe attributable to the promoted MA powder rearrangement due to Cu pools, which was formed during the sintering. Strch a rearrangement was well known to play an important role to densification at the liquid-phase sintering stage (2223). However, it must be noticed that the growth of the crystalline size is maintained to be constant even after the liquid-phase sintering (1 100°- 1300’) as shown in Figure 5. This may be explained as follows: after solid-state sintering, the W grains are totally surrounded by the Cu phase or Cu pools, so that W grains are coated with Cu phase. Thus, the W crystalline growth rate will be sharply decreased during the liquid-phase sintering stage, and the grain rearrangement occurs only at the liquid-phase sin&ring. This result is confirmed by the microstructural analysis which shows the isolated W particle in the sintered part at 1100” has still remained in the specimen sintered at 1300°, as shown in Figure 5. On the other hand, it has been reported that the grain growth in liquid-phase sir&ring is continued by the solution and reprecipitation due to the impurities effects induced by mechanical alloying process (24,25). However, in this study the grain size of W was kept to be constant with the regular size below 1 pm at liquid-phase sintering stage, and the rearrangement of W grain only occurs. We are, therefore,convinced that impurity effect is negligible in this liquid phase sit&ring. These results explain that the critical temperature of sintering is equal to the Cu melting temperature, because the W crystal size is hardly increased in the temperature range of over 1100”. CONCLUSIONS We fm:tly suggest the concepts of the nanosintering which explains the drastic crystalline growth of the nanostructured W-Cu alloys prepared by mechanical alloying during solid-stage sintering in the temperature range of 900’ to 1050”. Sintering of MA W-Cu alloys appears to be independent ‘ofCu content, and we found its origin in the thermal effect during heating-up stage. Sintering may be explained in terms of recovery and grain growth in the MA powders, as well as impurity activated sin&ring of W. After sintering, the Cu pools are formed outside MA powders. Due to these copper pools, therelative sintered density of more than 95% is achieved by the particle rearrangement during liquid phase-sintering. The critical temperature of sintering is equal to the Cu melting temperature, and the highest homogeneity and distribution of W and Cu phases is achieved by sintering at 1200’. ACKNOWLEDGMENT The authors acknowledge gratefully the financial support of the “Ministry of Education Research Fund for Advanced Materials in 1995.” REFERENCES 1.
German, R.M., Hens, K.F. and Johnson, J.L.,
2.
30, 205. Furukawa, N., Powder Metallurgy, 1990,33,213.
International
Journal of Powder Metallurgy, 1994,
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3. 4. 5. 6.
7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25.
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Moon, I.H., Proceedings of 1st Japan-Korea Workshop on Powder Injection Molding, Cheju, Korea, 1996. Lee, J.S. and Kim, T.H., Nanostructured Materials, 1995.6.691. Chongliang, Q., Enxi, W., Zhiqiang, E. and Yuhua, Z., Proceedings of the 13th International Plansee Seminar, ads., H. Bildstein and R. Eck, Metallwerk Plansee, Reutte, 1993, Vol. 1, p. 461. Aboud, T., Weiss, B.-Z. and Chaim, R., Nanostructured Materials, 1991, 6,405. Gaffet, E., Louison, C., Harmelin, M. and Faudot, F., Materials Science and Engineering1991, 134A, 1380. Kim, J.C., Ryu, S.S., Lee, H. and Moon, I.H. to be published in Journal of Korean Powder Metallurgy Institute. Benjamin, J.S. and Volin, J.E., Metallurgical Transactions, 1974,5, 1929. Johnson, J.L. andGerman,R.M.,Advancesin PowderMetallurgyandParticulateMaterials-1992, MPIF, Princeton, NJ, 1992, Vol. 3, p. 35. Bhalla, A.K. and Williams, J.D., Powder Metallurgy, 1976, 19,3 1. Yang, B. and German, R.M., Advances in Powder Metallurgy and Particulate Materials-1993, MPIF, Princeton, NJ, 1993, Vol. 2, p. 203. Kim, J.C. and Moon, I.H., to be submitted to Natwstructured Materials. Xiong, C.S., Xiong, Y.H., Zhu, H., Sun, T.F., Dong, E. and Liu, G.X., Nanostructured Materials, 1995,5,425. Drbohlav, 0. and Yavari, A.R., Acta Metallurgica et Materialia, 1995,43, 1799. Moon, I.H., Kim, J.H., Suk, M.J., Lee, K.M. and Lee, J.K., Journal of Refractory Metals and Hard Materials, 1992, 11,309. Moon, I.H., Kim, K.Y., Oh, S.T. and Suk, M.J., Journal ofAlloys and Compounds, 1993,201,129. Moon, I.H. and Lee, J.S., Powder Metallurgy International, 1977,9,23. Moon, I.H. and Lee, J.S., Powder Metallurgy, 1979,22,5. Miedema, A.R., Philips Tech. Review, 1976, 36,217. Panichkina, V.V., Sirotyuk, M.M. and Pilipovskii, Y.L., Poroshk. Metall., 1982,7, 1380. Kingery, W.D., Journal of Applied Physics, 1959,30,301. Cannon, H.S. and Lenel, F.V., Proceedings of the 1st International Plansee Seminar, eds., F. Benesovsky, Reutte, Austria, 1953, p. 106. Ahn, I.S., Kim, B.S. and Moon, I.H., Advances in Powder Metallurgy and Particulate Materials1995, MPIF, Princeton, NJ, 1995, Vol. 4, p. 209. Wittenauer, J. and Nieh, T.G., Tungsten and Tungsten Alloys Recent Advances, ed.s. A. Crowson and E.S. Chen, TMS, New Orleans, Louisiana, 1991, p. 21.
Peqpmon
Matmiek.Vol. 10. No. 2, pi. 283-290.1998 Ekevia ScienceLtd 8 1998AC&Mctdlurgi~ Inc. FvinkdintheUSA. All ri8ht.s ralavd 096%9773/98 $19.00 + .OO
PKISO9659773(98)00065-S
SINTERING OF NANOSTRUCTURED W-Cu ALLOYS PREPARED BY MECHANICAL ALLOYING Jin-Chun Kim and In-Hyung Moon Department of Materials Engineering, Hanyang University, Seoul 133-791, Korea (Accepted October 2,1997)
Absbwt - Nanwtructured (NS) powders with compositions corresponding to W2Owt%Cu and W-3Owt%Cu were prepared by mechanical alloying. The microstructure and grain size of as-milled and annealedpowders were analyzed by transmission electron microscopy. The compacted specimens were sintered at temperatures in the range lOOO”-13&I”, and then the microstructures of sintered parts were analyzed by scanning electron microscopy. Sintering of mechanicaily alloyed W-Cu alloysappears to be independent of Cu content, and may be explained in terms of recovery and grain growth in the mechanically alloyed powders as well as impurity activated simering of W. After sintering. Cu pools areformed outside the mechanically alloyed powders. A relative sintered density of more than 95% is obtained by particle rearrangement during liquid!-phase sintering, and the greatest homogeneity of W and Cu phases is achieved by sintering at i’2009 01998 Acta Metallurgica Inc.
INTRODUCTION Recendy, W-Cu alloys have been used for thermal managing and as microwave materials due to the high thermal conductivity of copper and the low thermal expansion coefficient of tungsten (l-3). The enhanced thermal properties of these alloys can be achieved by the unique combination of elemental phases. Research has been conducted on improving the homogeneity of component phases in these alloys. Among them, the use of nanostructured W-Cu powders prepared by coreduction of oxide powders (45) and mechanical alloying (MA) of elemental powders (6,;‘) has been widely studied. We investigated the characteristics of nanostructured W-Cu alloys prepared by MA (8). In this experiment, it was found that the grain growth of the MA powders significantly occurred at solid-state sintering temperatures below 1083O. The internal structure of MA powders after sintering was similar to the conventional one of W-Cu alloys processed by liquid-phase s&ring. The enhanced sinterability of MA W-Cu alloys was also obtained during solid-state sintering. After sintcring, a high sintered density and a homogeneity of W-Cu alloys were obtained. In the present work, the sintering behavior of MA W-Cu alloys is described and a possible explanation for the observations is suggested. 283
J-C KIM AND1-H kGN
EXPERIMENTAL
Elemental W (99.9% purity) and Cu (99.5% purity) powders with particle size of 4.28 l.trn and 50.42 p.m, respectively, were used in this experiment. Composite powders with compositions corresponding to the W-20wt%Cu and W-30wt%Cu were prepared by mechanical alloying. The MA was performed in an attritor of 750 cm3 capacity using stainless steel balls with 0.48 cm diameter. The weight ratio of ball to powder was maintained to60: 1. The microstructure and grain size of as-milled powders and annealed powders at 900” were analyzed by transmission electron microscopy (TEM, model:2000-EX II, JEOL) by attaching the powders on the 300 mesh grid. The bulk specimens of MA powders were prepared by cold pressing into a disk shape with a green density of 40&2% of theoretical density. The compacted MAW-3OCu specimens were sintered in the temperature range looO”-1300” for 1 hr in HZatm. The sintering of the MAW-2OCu compacts was carried out at 1000° and 1050’ for 0 minute in order to investigate the microstructural change of specimens during the heat-up stage( 0 min. means that the holding time at the sintering temperature is zero). The microstructures of sintered parts were analyzed by scanning electron microscopy (SEM).
RESULTS
In mechanical alloying, the development of particle shape and size reaches steady-state through several different stages of processing, as suggested by Benjamin (9). From the observation of morphological development and change of x-ray diffraction patterns with MAtimes, the milling time of the steady-state condition in this experiment was 100 hrs. In that case, the MA powders were characterized by three dimensional equiaxed shape, and its size was 3-5 l.trn.
Figure 1. The TEM micrographs of the mechanically alloyed W-30wt%Cu powder for 100 hr; bright field TEM image (a) and selected area diffraction (SAD) pattern (b).
NANOSINTERING OF W-Cu ALLOYSPREPAREO BY MA
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Figure 1 shows the bright field TEM image and selected area diffraction (SAD) pattern of MA W-30& powders at the steady-state condition. As shown in this figure, the grain size of MA W-Cu powders is 20 to 30 nm, and the SAD pattern of matrix with sharp diffracted rings indicates that the W and Cu remain in the crystalline state. In addition, the irregular grain morphology of the as-milled powder was observed and W phases are not distinguishable from the Cu phases. Figure 2 shows the SEM and TEM microstructures of the MA W-3001 powders annealed at 900” in Hz atmosphere. As shown in these figures, the grain size of MApowders (in Figure 2(b)) with annealing for 5 hrs increases from 20-30 nm of as-milled powders to about 130 nm. Now, we are able to distinguish W from Cu in the annealed MA powders, and it shows that the W grains undergo shape change from irregular shape to spheroid. In addition, some of the W-grains are coalesced or even sintered (marked with a circle). They exhibited similar microstructural properties fa.bricated by the conventional liquid-phase sintering in the W-Cu system using microsized powers (10.11). This coalescence of W/W grains is clearly visible in the TEM microstructure (Figure 2(c):). In this figure, a neck formation due to W-W sintering is evident (arrow A), and also W grain is bounded by the Cu phase (arrow B) and surrounded by the Cu layer. However, the external morphology of MA powder (Figure 2(a)) is almost identical to that of initial MA powder regardless of different thermal treatment. Figure 3 shows the SEM microstructures of MA W-3OCu compact sintered at 1000” for 1 hr. The W-grain size of the MA powders increases to about 200 nm at this temperature. The interior microstructure of the MA W-Cu powders observed with SEM exhibits a typical microstructure of fully densified W-Cu alloys by liquid phase sintering or Cu infiltration method (12). It is also observed that the Cu rich phase or pools indicated by the arrow covers the exterior of MA powders. The morphology of MA powders in the sintered part (Figure 3(a)) is similar to that of initial MA
Figure 2. The SEM micrographs of powder (a) and internal structure (b) and bright field TEM micrograph (c) of the MA W-30wt%Cu powders annealed at 900” in H2 atm. for 5 hr.
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powder and annealed powder at 900°, and the MA powders are aggregated by the Cu pools. The details on the formation of Cu pools at solid-stage sintering will be discussed in the separate paper (13). As shown in Figures 2 and 3, the grain size of the MA powders drastically increases from 20-30 nm to 200 nm at the solid-stage sintering stage and theinternal microstructure of MA powder which was formed from the coalescence or sintering of W grains is similar to that of typical liquidphase sintering of the W-Cu alloys. If the sintering of the MA W-Cu powders takes place, the Wgrain size of the MA powders drastically increases from the initially fine-grained as-milled powders withoutchanging thepowdershapeandtheintemai microstmctureofMApowderis very similar to that of the liquid-phase sintering, and notably the Cu pools are formed outside MA
Figure 3. The SEM micrographs of the MA W-3Owt%Cu sintered at 1000’ for 1 hr in Hz atm.; Cu pool area (a) and interior area of powder (b).
Figure 4. The SEM micrographs of the MA W-20wtlCu sintered for 0 min in HZ atm. at various temperatures; 1000° (a) and 1050’ (b).
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Figure ,4(a) and (b) exhibit the microstructure of MA W-20wt%Cu specimens sintered at 1000” and 10!jO”for0 min in HZ atm,respectively. As shown in these figures, sintering also occurs, and its shucture is well matched to that of the MA W-3OCu specimen. Based on this observation, we suggest thlat the nanosintering of MA W-Cu powders may be independent of Cu content, and we find its origin in the thermal effects of the heating-up stage. Figure 5 shows SEM micrographs of MA W-3Owt%Cu specimens sintered at 1100” (a), 1150” (b), 12OOO(c)and13OO”(d)for 1 hrinH2atm. TheW-grainsizeof8OOto lOOOnm,asshown in these figums, is much larger compared to that of the specimens sintered at solid-state sintering, and the mixing state of W and Cu phase is more homogeneous with increasing sintering temperature. The sintered density of all specimens is higher than the relative density of 95%, and there is no observable porosity. The MA powder shape in the specimen sintered at 1100’ is not changed fmm the initial morphology of MA powders, but the W grains are totally surrounded by the Cu liquid phase. At 1150’ (in Figure 5(b)), the shape of agglomerated W grains begins to change to the homogeneous state without further growth of the W grains. In the case of the specimen sir&red at 1200’ (Figure 5(c)), it shows that each W grain in the MA powder is rearranged, alnd the highest homogenous state of W and Cu phases can be attained by sintering. Finally, the microstructure of the specimen sintered at 1300” is observed to be analogous to that of the liquid-phase sintering investigated by others, as shown in Figure 5(d).
DISCUSSION 1. Solid-stateSintering In general, the grain growth of the W-Cu alloys during the solid-state sintering is negligible because W and Cu are completely immiscible. However, as described above, the grain size of MA powders drasticallyincreases during the solid-state sintering and its inner microstructure gradually changes to tlhat of conventional liquid-phase sintering. This phenomenon may be explained in terms of following mechanisms. Firstly, it may be derived from the recovery and grain growth due to internal strain energy, which is caused by the repeated cold welding and fracture during mechanical alloying process. The stored energy of the MA powders will be released in the form of external thermal effect during heating (14,15). Consequently, it can be assumed that the unstable W and Cu grains and lattice deformations in the MA powders transform into the stable state with increasing temperature, i.e., the recovery and recrystallization occur during the heating, and then the grain growth occurs in the MA powders. The second possibility is the activated sintering of W due to the Fe impurity, which is inherent in the MA powders from the ball media and attritor impeller during the MA process. The effect of impurity addition such as Fe, Ni and Co etc., which causes the activated sintering of W, was well known to enhance the sintering behavior of W (16,17). Third,. the activated liquid phase sintering of W-Cu due to impurity must be considered (10,18,19). The addition of the activators which are soluble in the liquid phase enhances densification and grain growth through solution and reprecipitation. However, it should be carefully colnsidered that the impurity effect on W-Cu alloys only acts at the liquid-phase sintering and it shouldbe accompanied with the continuous grain growth in the liquid phase sintering stage.
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Figure 5. The micrographs observed by the MA W-30wt%Cu sintered for 1 hr at various temperatures: 1100’ (a), 1150 (b), 1200 (c) and 1300’ (d). In our study, the growth of W grain is solely observed during the solid-state sintering, and no significant grain growth occurs in the liquid-phase sintering. Finally, we should consider the interreaction of the Wand the Cu constituents. Under the equilibrium condition, W and Cu are completely immiscible due to the high positive heat of mixing; 35SKJ/mol(20). But, Panichkina (21) reported that the considerable grain growth of W particle was observed with increasing W solubility in Cu, which resulted from the use of W particle of sub-micron size. However, this mechanism cannot be directly applied to the W-grain growth in the MA powders at the solid-state sintering because his experiment was carried out only in the range of the liquid-phase sintering. Besides the above mechanisms, the solubility increase of Cu in W due to the MA as well as the decrease of Cu melting temperature due to the Gibbs-Thomson effect should be taken into account. Among them, the mechanism of recovery followed by grain growth of the MA powder and the impurity activated sintering of W seem to be the case in this experiment. No matter what mechanisms operate, however, the drastic grain growth in the NS W-Cu powders prepared by mechanical alloying, i.e., nanosintering occurred, and the Cu pools were formed in the external of the MA powders.
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2. Microstructure Changes of Liquid Phase Sintering Stage
In this experiment, the relative sintered density of the nanostructured W-Cu alloys increased drastically above 95% in the solid-stage sintering stage at 1050’ and the high sintered density of 98% wasobtainedat liquid-phase sintering temperatures (at 1100°). This result maybe attributable to the promoted MA powder rearrangement due to Cu pools, which was formed during the sintering. Strch a rearrangement was well known to play an important role to densification at the liquid-phase sintering stage (2223). However, it must be noticed that the growth of the crystalline size is maintained to be constant even after the liquid-phase sintering (1 100°- 1300’) as shown in Figure 5. This may be explained as follows: after solid-state sintering, the W grains are totally surrounded by the Cu phase or Cu pools, so that W grains are coated with Cu phase. Thus, the W crystalline growth rate will be sharply decreased during the liquid-phase sintering stage, and the grain rearrangement occurs only at the liquid-phase sin&ring. This result is confirmed by the microstructural analysis which shows the isolated W particle in the sintered part at 1100” has still remained in the specimen sintered at 1300°, as shown in Figure 5. On the other hand, it has been reported that the grain growth in liquid-phase sir&ring is continued by the solution and reprecipitation due to the impurities effects induced by mechanical alloying process (24,25). However, in this study the grain size of W was kept to be constant with the regular size below 1 pm at liquid-phase sintering stage, and the rearrangement of W grain only occurs. We are, therefore,convinced that impurity effect is negligible in this liquid phase sit&ring. These results explain that the critical temperature of sintering is equal to the Cu melting temperature, because the W crystal size is hardly increased in the temperature range of over 1100”. CONCLUSIONS We fm:tly suggest the concepts of the nanosintering which explains the drastic crystalline growth of the nanostructured W-Cu alloys prepared by mechanical alloying during solid-stage sintering in the temperature range of 900’ to 1050”. Sintering of MA W-Cu alloys appears to be independent ‘ofCu content, and we found its origin in the thermal effect during heating-up stage. Sintering may be explained in terms of recovery and grain growth in the MA powders, as well as impurity activated sin&ring of W. After sintering, the Cu pools are formed outside MA powders. Due to these copper pools, therelative sintered density of more than 95% is achieved by the particle rearrangement during liquid phase-sintering. The critical temperature of sintering is equal to the Cu melting temperature, and the highest homogeneity and distribution of W and Cu phases is achieved by sintering at 1200’. ACKNOWLEDGMENT The authors acknowledge gratefully the financial support of the “Ministry of Education Research Fund for Advanced Materials in 1995.” REFERENCES 1.
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