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Microstructure analysis of Co–Cu alloys undercooled prior to solidification
Materials Science and Engineering A 375–377 (2004) 1101–1104
Microstructure analysis of Co–Cu alloys undercooled prior to solidification X.Y. Lu a,b , C.D. Cao b , M. Kolbe a,∗ , B. Wei b , D.M. Herlach a b
a Institute of Space Simulation, German Aerospace Center (DLR), D-51170 Cologne, Germany Department of Applied Physics, Northwestern Polytechnical University, P.O. Box 624, Xian 710072, PR China
Abstract Co41.8 Cu58.2 and Co16 Cu84 alloys were investigated by electromagnetic levitation (EML) and drop tube (DT) processing. The experiments show that both methods allow deep undercooling of the alloys into the miscibility gap. By electromagnetic levitation, in Co41.8 Cu58.2 alloy the distribution of Co-rich phases was strongly influenced by electromagnetic stirring, while in Co16 Cu84 alloy the Co-rich phases exhibit rounder shape. By drop tube experiment, the disturbance of the electromagnetic stirring is reduced, the Co-rich phases distribute homogeneously in the matrix phase. © 2003 Elsevier B.V. All rights reserved. Keywords: Rapid solidification; Undercooling; Phase separation; Microstructure
1. Introduction A special phenomenon was found in peritectic systems like Co–Cu or Fe–Cu, which have nearly flat liquidus line and retrograde solidus line and hence display a thermodynamic tendency towards the formation of a miscibility gap in undercooled liquid state [1]. The Co–Cu alloy melt will separate into Co-rich L1-phase and Cu-rich L2-phase when it is undercooled into the metastable miscibility gap. In 1958, Nakagawa [1] determined the magnetic susceptibility of undercooled samples and found the miscibility gap in Co–Cu and Fe–Cu systems for the first time. Since then the research interest has been aroused in the phase separation effect in both systems. For Co–Cu system, Munitz and coworkers [2–5] had done much work by electromagnetic levitation (EML), drop tube (DT) and electron beam surface melting on the undercooling of bulk samples and obtained the miscibility gap of this system. Yamauchi et al. [6] investigated the undercooling and its effect on solidification structure by glass fluxing methods. Sun et al. [7] studied the supercooling and liquid phase separation by using isothermal soaking at different supercooling rates. But there is still some research space in this system in the mechanism of liquid phase
∗
Corresponding author. E-mail address: [email protected] (M. Kolbe).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2003.10.106
separation and the solidification behaviour of the separated phases. In this work, in order to investigate the effect of undercooling, cooling rate and composition on liquid phase separation, methods such as electromagnetic levitation and drop tube are used to study the microstructure evolution of Co41.8 Cu58.2 and Co16 Cu84 alloys in the metastable miscibility gap. The different alloys will exhibit the influence of the volume fraction of the separated minority phase on the coagulation of this phase.
2. Experimental procedure In the present work, microstructure evolution of Co41.8 Cu58.2 and Co16 Cu84 alloys under different solidification conditions were investigated. Samples with weight around 1 g were prepared from 99.999% pure Cu and 99.998% pure Co by arc-melting under Ar atmosphere. Both alloys were processed by the electromagnetic levitation facility and in a 8 m-drop tube. The details of the EML facility and drop tube are based on the facilities described elsewhere [8,9]. After experiments, the samples were analysed according to standard metallographic procedures. The solidification microstructures were investigated by means of scanning electron microscope (SEM) LEO 1530 VP from LEO Elektronenmkroskopie with detectors for backscattered electrons (BSE). The local chemical composition of minor and major
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phases of the samples has been measured with an Oxford Instruments energy dispersive X-ray analysis (EDX) system.
3. Results and discussions The composition range of the Co41.8 Cu58.2 and Co16 Cu84 alloys are shown in the Co–Cu binary phase diagram [10] of Fig. 1. The solid circles are the precisely measured miscibility gap by Cao et al. [11,12], the dashed line is the graphical fitting curve of the composition dependence of the binodal temperature. It can be seen that when the liquid melt enters the metastable region after sufficient undercooling, the melt separates into a Co-rich L1-phase and a Cu-rich L2-phase. The compositions of both phases will follow the boundary of the miscibility gap during further cooling. 3.1. Microstructures of Co41.8 Cu58.2 alloy The liquidus temperature of the Co41.8 Cu58.2 alloy is 1651 K, the interval between the liquidus and binodal temperature is about 105 K [11]. If the sample solidified in an equilibrium way, the microstructure of the sample at room temperature would be 39% ε-Co phase plus 61% Cu phase in volume fraction. The maximum undercooling obtained by EML is 207 K, which is indicated in Fig. 1 by a solid square. Therefore the expected microstructure should reveal the separation into L1- and L2-phase before solidification. Fig. 2 shows microstructures of this alloy solidified at different undercoolings. In Fig. 2a, the microstructure is dendrites when the sample was undercooled only slightly to 17 K. However, there are different morphologies of the dendrites. In some sites the dendrites keep their intact shapes, but in some region there are some blocks of dendrites, this results from the remelting of the dendrite during recalescence when the release of the heat of crystallization leads to a rapid increase 1800 Exp. Fit.
Co41.8 Cu58.2 L
1700
Temperature (K)
Co16 Cu84 L+ α(Co)
α (Co)
1600
*
∆ T =105K ∆T =116K *
L1+L2
1500
∆ T=207K ∆T=193K
1385K
1400
Cu 1300
0
20
40
60
80
at.% Cu Fig. 1. Phase diagram of Co–Cu binary alloys.
100
Fig. 2. Microstructure of Co41.8 Cu58.2 alloys at different undercoolings by EML. (a) T = 17 K and (b) T = 207 K.
of the temperature of the levitation processed sample. But either in the entire dendrite or in the fragmentation of dendrite, the average composition of the dendrite is 83 at.% Co, and the matrix contains only 14 at.% Co. Fig. 2b shows the morphology of phase separation. The sample is undercooled to 207 K that corresponds to a cooling of 102 K into the miscibility gap. Fig. 2b only shows a part of the sample. The black phase corresponds to Co-rich L1-phase, the matrix is formed by Cu-rich L2-phase. The region of the L1-phase is distorted due to the electromagnetic stirring of the melt. The coalescence of the separated phase is strongly developed with the reduction of the interface energy as its driving force. As shown in Fig. 2b, some Co-rich droplets meet together and lose their original shapes and conform somewhat to the shape of the droplet meeting with them. It also can be seen that some ␣-Co dendrites grew from the periphery of the Co-rich L1-phase, indicating that the Co-rich L1-phase could act as their nuclei. The composition of the Co-rich phase is determined as 84 at.% Co, while the matrix contains about 9 at.% Co. Fig. 3 shows the morphology of the sample solidified in drop tube. The liquid separated into L1- and L2-phase before solidification. The size of the droplet is 180 m. It can be observed that there is a size transformation range of the Co-rich spheres. In the middle of the droplet, the Co-rich spheres are finer, the typical size of the spheres is about 1.5 m, while in the edge of the droplet the typical size is about 3 m.
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regime. Since dissolution of demixed particles increases with temperature the particles in the middle of the droplets are smaller than in the outer regime. Therefore, the final microstructure has a finer structure in the middle than in the edge of the droplet. 3.2. Microstructures of Co16 Cu84 alloy
Fig. 3. Microstructure of Co41.8 Cu58.2 alloy by drop tube.
Several mechanisms have been proposed to explain the evolution of size distribution of the dispersion phase, including wetting behavior, Ostwald ripening, and coarsening of droplets by collision. Ostwald ripening is a mechanism, which explains droplets coalescence by solute diffusion between them. The solute solubility depends on the curvature of a droplet, the smaller the curvature, i.e. the larger the radius, the smaller the solubility [13]. This diffusion dependent coarsening mechanism plays a dominant role when the dispersion phase has small diameter. With the increase of droplet radius and the decrease of temperature, the solid–liquid interface tension increases, whereas the solubility decreases, which results in the weakening of Ostwald ripening. In such a case, collision mechanism is likely to dominate the growth of dispersion phase, which involves relative motion, collision and subsequent coalescence of the droplets. Collision of the droplets results mainly from Stokes motion and Marangoni migration. Marangoni migration is caused by surface tension gradient of the dispersion droplets in the matrix phase, which is related to the gradients of concentration and temperature [14]. The droplet tends to move in the direction of surface tension decreasing. According to the measurement of Eichel and Egry [15], the surface tension of Co–Cu alloy decrease with the increase of temperature. If the coagulation of Co-rich phases is controlled by Marangoni convection, the separated Co-rich spheres should move from the surface to the center of the droplet due to the small temperature gradient inside the droplet. But it could be found in Fig. 3 that the separated Co-rich spheres moved from the center to the edge of the droplet, this might be due to the temperature gradient induced solute diffusion. If two separated spheres have same size but with different temperature, the solute diffusion in the higher temperature one is much faster than in the lower temperature one. The rapid solidification of the sample after deep undercooling leads to freeze in the instantaneous state of demixing and a rapid release of the heat of crystallization at the solid–liquid interface. Since this heat production is more rapidly removed at the boundary region near the surface of the sample compared to the volume, the interior of the sample heats up to higher temperatures than the surface
The liquidus temperature of the Co16 Cu84 alloy is 1583 K, the interval between the liquidus and binodal temperature amounts to 116 K [11]. Under conditions of equilibrium solidification, the microstructure of the sample at room temperature would be 14.5% ε-Co phase plus 85.5% Cu phase in volume fraction. A maximum undercooling of 193 K is obtained by EML, which is also shown in Fig. 1 by a solid square. Such an undercooling corresponds to a temperature being 77 K below the binodal temperature of this alloy composition. Fig. 4 shows microstructures of this alloy at different undercoolings. According to Fig. 4a, the sample solidified in dendrite morphology at undercooling of 65 K. The average composition of the dendrite is 78 at.% Co, while the matrix contains only 5 at.% Co. This is much lower than that in Co41.8 Cu58.2 alloy. Also the size and the amount of the dendrites are rather lower than that in Co41.8 Cu58.2 alloy. Fig. 4b shows the morphology of phase separation. It is a little different from that of Co41.8 Cu58.2 alloy. Here most of Co-rich phases solidified as round spheres. The average composition
Fig. 4. Microstructure of Co16 Cu84 alloy at different undercoolings by EML. (a) T = 65 K and (b) T = 193 K.
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undercoolings were achieved down to 207 and 193 K, respectively, the alloy melts separated into two liquid phases before solidification. For the Co41.8 Cu58.2 alloy the microstructure was strongly influenced by the electromagnetic stirring in EML experiment, while in Co16 Cu84 alloy the Co-rich phases exhibit rounder shape. In drop tube processing, the Co-rich phases disperse more homogeneously than in EML. This is due to the reduced gravity and no strong electromagnetic stirring during free fall and solidification. Acknowledgements Fig. 5. Microstructure of Co16 Cu84 alloy by drop tube.
of them is 83 at.% Co, the matrix contains 4 at.% Co. This indicates that the composition plays an important role during the liquid phase separation. In Co16 Cu84 alloy, the volume fraction of separated Co-rich L1-phase is small and the coagulation of L1-phase is not strongly distorted by electromagnetic stirring. Under deep undercooling, when liquid phase separation occurred, Co41.8 Cu58.2 and Co16 Cu84 alloys contain different amount of Co-rich L1-phase. The higher the volume fraction of the Co-rich L1-phase, the larger the opportunity of them to coagulate. This explains the larger size of Co-rich phase obtained in Co41.8 Cu58.2 alloy during EML experiment compared with Co16 Cu84 alloy. Fig. 5 shows the sample of Co16 Cu84 alloy solidified in drop tube. The size of the droplet is 150 m. The Co-rich L1-phase solidified as sphere and nearly dispersed homogeneously in the entire droplets. The largest diameter of the L1-phase is about 3.3 m. Comparing the samples solidified in EML and DT processing, the electromagnetic stirring and density difference induced stokes motion influences the coalescence of the L1-phase in EML procedure. While in drop tube experiment, both effects are suppressed, the minor phase distributes homogeneously in the matrix phase. 4. Conclusion Both Co41.8 Cu58.2 and Co16 Cu84 alloys were treated by electromagnetic levitation and drop tube processing. Deep
This work is financially supported by the National Natural Science Foundation of China (Grant No. 50221101). One of the authors (X.Y. Lu) is grateful to BMBF for a visiting scholarship. The authors gratefully acknowledge Prof. L. Ratke, Ms. X.R. Liu, Dr. J.R. Gao, Dr. T. Volkmann and Dr. O. Funke for useful discussions and their help with experiments. References [1] Y. Nakagawa, Acta Metall. 6 (1958) 704. [2] S.P. Elder, A. Munitz, G.J. Abbaschian, Mater. Sci. Forum 50 (1989) 137. [3] A. Munitz, R. Abbaschian, Metall. Mater. Trans. 27A (1996) 4049. [4] A. Munitz, S.P. Elder, R. Abbaschian, Metall. Mater. Trans. 23A (1992) 1817. [5] A. Munitz, R. Abbaschian, J. Mater. Sci. 33 (1998) 3639. [6] I. Yamauchi, N. Ueno, M. Shimaoka, I. Ohnaka, J. Mater. Sci. 33 (1998) 371. [7] Z. Sun, X. Song, Z. Hu, S. Yang, G. Liang, J. Sun, J. Alloys Comp. 319 (2001) 266. [8] F. Gillessen, Forschungsbericht DFVLR, DLR-FB 89-32, Köln, 1989. [9] D.M. Herlach, R. Willnecker, F. Gillessen, in: Proceedings of the 5th European Symposium on Materials Science under Microgravity, ESA SP-2 22, European Space Agency, Noordwijk, The Netherlands, 1985, p. 399. [10] T. Nishizawa, K. Ishida, Bull. Alloy Phase Diagrams 5 (1984) 161. [11] C.D. Cao, G.P. Görler, D.M. Herlach, B. Wei, Mater. Sci. Eng. A 325 (2002) 503. [12] C.D. Cao, T. Letzig, G.P. Görler, D.M. Herlach, J. Alloys Comp. 325 (2001) 113. [13] J.R. Rogers, R.H. Davis, Metall. Trans. 21A (1990) 59. [14] L. Ratke, S. Diefenbach, Mater. Sci. Eng. R 15 (1995) 263. [15] R. Eichel, I. Egry, Z. Metallkd. 90 (1999) 371.
Microstructure analysis of Co–Cu alloys undercooled prior to solidification X.Y. Lu a,b , C.D. Cao b , M. Kolbe a,∗ , B. Wei b , D.M. Herlach a b
a Institute of Space Simulation, German Aerospace Center (DLR), D-51170 Cologne, Germany Department of Applied Physics, Northwestern Polytechnical University, P.O. Box 624, Xian 710072, PR China
Abstract Co41.8 Cu58.2 and Co16 Cu84 alloys were investigated by electromagnetic levitation (EML) and drop tube (DT) processing. The experiments show that both methods allow deep undercooling of the alloys into the miscibility gap. By electromagnetic levitation, in Co41.8 Cu58.2 alloy the distribution of Co-rich phases was strongly influenced by electromagnetic stirring, while in Co16 Cu84 alloy the Co-rich phases exhibit rounder shape. By drop tube experiment, the disturbance of the electromagnetic stirring is reduced, the Co-rich phases distribute homogeneously in the matrix phase. © 2003 Elsevier B.V. All rights reserved. Keywords: Rapid solidification; Undercooling; Phase separation; Microstructure
1. Introduction A special phenomenon was found in peritectic systems like Co–Cu or Fe–Cu, which have nearly flat liquidus line and retrograde solidus line and hence display a thermodynamic tendency towards the formation of a miscibility gap in undercooled liquid state [1]. The Co–Cu alloy melt will separate into Co-rich L1-phase and Cu-rich L2-phase when it is undercooled into the metastable miscibility gap. In 1958, Nakagawa [1] determined the magnetic susceptibility of undercooled samples and found the miscibility gap in Co–Cu and Fe–Cu systems for the first time. Since then the research interest has been aroused in the phase separation effect in both systems. For Co–Cu system, Munitz and coworkers [2–5] had done much work by electromagnetic levitation (EML), drop tube (DT) and electron beam surface melting on the undercooling of bulk samples and obtained the miscibility gap of this system. Yamauchi et al. [6] investigated the undercooling and its effect on solidification structure by glass fluxing methods. Sun et al. [7] studied the supercooling and liquid phase separation by using isothermal soaking at different supercooling rates. But there is still some research space in this system in the mechanism of liquid phase
∗
Corresponding author. E-mail address: [email protected] (M. Kolbe).
0921-5093/$ – see front matter © 2003 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2003.10.106
separation and the solidification behaviour of the separated phases. In this work, in order to investigate the effect of undercooling, cooling rate and composition on liquid phase separation, methods such as electromagnetic levitation and drop tube are used to study the microstructure evolution of Co41.8 Cu58.2 and Co16 Cu84 alloys in the metastable miscibility gap. The different alloys will exhibit the influence of the volume fraction of the separated minority phase on the coagulation of this phase.
2. Experimental procedure In the present work, microstructure evolution of Co41.8 Cu58.2 and Co16 Cu84 alloys under different solidification conditions were investigated. Samples with weight around 1 g were prepared from 99.999% pure Cu and 99.998% pure Co by arc-melting under Ar atmosphere. Both alloys were processed by the electromagnetic levitation facility and in a 8 m-drop tube. The details of the EML facility and drop tube are based on the facilities described elsewhere [8,9]. After experiments, the samples were analysed according to standard metallographic procedures. The solidification microstructures were investigated by means of scanning electron microscope (SEM) LEO 1530 VP from LEO Elektronenmkroskopie with detectors for backscattered electrons (BSE). The local chemical composition of minor and major
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phases of the samples has been measured with an Oxford Instruments energy dispersive X-ray analysis (EDX) system.
3. Results and discussions The composition range of the Co41.8 Cu58.2 and Co16 Cu84 alloys are shown in the Co–Cu binary phase diagram [10] of Fig. 1. The solid circles are the precisely measured miscibility gap by Cao et al. [11,12], the dashed line is the graphical fitting curve of the composition dependence of the binodal temperature. It can be seen that when the liquid melt enters the metastable region after sufficient undercooling, the melt separates into a Co-rich L1-phase and a Cu-rich L2-phase. The compositions of both phases will follow the boundary of the miscibility gap during further cooling. 3.1. Microstructures of Co41.8 Cu58.2 alloy The liquidus temperature of the Co41.8 Cu58.2 alloy is 1651 K, the interval between the liquidus and binodal temperature is about 105 K [11]. If the sample solidified in an equilibrium way, the microstructure of the sample at room temperature would be 39% ε-Co phase plus 61% Cu phase in volume fraction. The maximum undercooling obtained by EML is 207 K, which is indicated in Fig. 1 by a solid square. Therefore the expected microstructure should reveal the separation into L1- and L2-phase before solidification. Fig. 2 shows microstructures of this alloy solidified at different undercoolings. In Fig. 2a, the microstructure is dendrites when the sample was undercooled only slightly to 17 K. However, there are different morphologies of the dendrites. In some sites the dendrites keep their intact shapes, but in some region there are some blocks of dendrites, this results from the remelting of the dendrite during recalescence when the release of the heat of crystallization leads to a rapid increase 1800 Exp. Fit.
Co41.8 Cu58.2 L
1700
Temperature (K)
Co16 Cu84 L+ α(Co)
α (Co)
1600
*
∆ T =105K ∆T =116K *
L1+L2
1500
∆ T=207K ∆T=193K
1385K
1400
Cu 1300
0
20
40
60
80
at.% Cu Fig. 1. Phase diagram of Co–Cu binary alloys.
100
Fig. 2. Microstructure of Co41.8 Cu58.2 alloys at different undercoolings by EML. (a) T = 17 K and (b) T = 207 K.
of the temperature of the levitation processed sample. But either in the entire dendrite or in the fragmentation of dendrite, the average composition of the dendrite is 83 at.% Co, and the matrix contains only 14 at.% Co. Fig. 2b shows the morphology of phase separation. The sample is undercooled to 207 K that corresponds to a cooling of 102 K into the miscibility gap. Fig. 2b only shows a part of the sample. The black phase corresponds to Co-rich L1-phase, the matrix is formed by Cu-rich L2-phase. The region of the L1-phase is distorted due to the electromagnetic stirring of the melt. The coalescence of the separated phase is strongly developed with the reduction of the interface energy as its driving force. As shown in Fig. 2b, some Co-rich droplets meet together and lose their original shapes and conform somewhat to the shape of the droplet meeting with them. It also can be seen that some ␣-Co dendrites grew from the periphery of the Co-rich L1-phase, indicating that the Co-rich L1-phase could act as their nuclei. The composition of the Co-rich phase is determined as 84 at.% Co, while the matrix contains about 9 at.% Co. Fig. 3 shows the morphology of the sample solidified in drop tube. The liquid separated into L1- and L2-phase before solidification. The size of the droplet is 180 m. It can be observed that there is a size transformation range of the Co-rich spheres. In the middle of the droplet, the Co-rich spheres are finer, the typical size of the spheres is about 1.5 m, while in the edge of the droplet the typical size is about 3 m.
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regime. Since dissolution of demixed particles increases with temperature the particles in the middle of the droplets are smaller than in the outer regime. Therefore, the final microstructure has a finer structure in the middle than in the edge of the droplet. 3.2. Microstructures of Co16 Cu84 alloy
Fig. 3. Microstructure of Co41.8 Cu58.2 alloy by drop tube.
Several mechanisms have been proposed to explain the evolution of size distribution of the dispersion phase, including wetting behavior, Ostwald ripening, and coarsening of droplets by collision. Ostwald ripening is a mechanism, which explains droplets coalescence by solute diffusion between them. The solute solubility depends on the curvature of a droplet, the smaller the curvature, i.e. the larger the radius, the smaller the solubility [13]. This diffusion dependent coarsening mechanism plays a dominant role when the dispersion phase has small diameter. With the increase of droplet radius and the decrease of temperature, the solid–liquid interface tension increases, whereas the solubility decreases, which results in the weakening of Ostwald ripening. In such a case, collision mechanism is likely to dominate the growth of dispersion phase, which involves relative motion, collision and subsequent coalescence of the droplets. Collision of the droplets results mainly from Stokes motion and Marangoni migration. Marangoni migration is caused by surface tension gradient of the dispersion droplets in the matrix phase, which is related to the gradients of concentration and temperature [14]. The droplet tends to move in the direction of surface tension decreasing. According to the measurement of Eichel and Egry [15], the surface tension of Co–Cu alloy decrease with the increase of temperature. If the coagulation of Co-rich phases is controlled by Marangoni convection, the separated Co-rich spheres should move from the surface to the center of the droplet due to the small temperature gradient inside the droplet. But it could be found in Fig. 3 that the separated Co-rich spheres moved from the center to the edge of the droplet, this might be due to the temperature gradient induced solute diffusion. If two separated spheres have same size but with different temperature, the solute diffusion in the higher temperature one is much faster than in the lower temperature one. The rapid solidification of the sample after deep undercooling leads to freeze in the instantaneous state of demixing and a rapid release of the heat of crystallization at the solid–liquid interface. Since this heat production is more rapidly removed at the boundary region near the surface of the sample compared to the volume, the interior of the sample heats up to higher temperatures than the surface
The liquidus temperature of the Co16 Cu84 alloy is 1583 K, the interval between the liquidus and binodal temperature amounts to 116 K [11]. Under conditions of equilibrium solidification, the microstructure of the sample at room temperature would be 14.5% ε-Co phase plus 85.5% Cu phase in volume fraction. A maximum undercooling of 193 K is obtained by EML, which is also shown in Fig. 1 by a solid square. Such an undercooling corresponds to a temperature being 77 K below the binodal temperature of this alloy composition. Fig. 4 shows microstructures of this alloy at different undercoolings. According to Fig. 4a, the sample solidified in dendrite morphology at undercooling of 65 K. The average composition of the dendrite is 78 at.% Co, while the matrix contains only 5 at.% Co. This is much lower than that in Co41.8 Cu58.2 alloy. Also the size and the amount of the dendrites are rather lower than that in Co41.8 Cu58.2 alloy. Fig. 4b shows the morphology of phase separation. It is a little different from that of Co41.8 Cu58.2 alloy. Here most of Co-rich phases solidified as round spheres. The average composition
Fig. 4. Microstructure of Co16 Cu84 alloy at different undercoolings by EML. (a) T = 65 K and (b) T = 193 K.
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undercoolings were achieved down to 207 and 193 K, respectively, the alloy melts separated into two liquid phases before solidification. For the Co41.8 Cu58.2 alloy the microstructure was strongly influenced by the electromagnetic stirring in EML experiment, while in Co16 Cu84 alloy the Co-rich phases exhibit rounder shape. In drop tube processing, the Co-rich phases disperse more homogeneously than in EML. This is due to the reduced gravity and no strong electromagnetic stirring during free fall and solidification. Acknowledgements Fig. 5. Microstructure of Co16 Cu84 alloy by drop tube.
of them is 83 at.% Co, the matrix contains 4 at.% Co. This indicates that the composition plays an important role during the liquid phase separation. In Co16 Cu84 alloy, the volume fraction of separated Co-rich L1-phase is small and the coagulation of L1-phase is not strongly distorted by electromagnetic stirring. Under deep undercooling, when liquid phase separation occurred, Co41.8 Cu58.2 and Co16 Cu84 alloys contain different amount of Co-rich L1-phase. The higher the volume fraction of the Co-rich L1-phase, the larger the opportunity of them to coagulate. This explains the larger size of Co-rich phase obtained in Co41.8 Cu58.2 alloy during EML experiment compared with Co16 Cu84 alloy. Fig. 5 shows the sample of Co16 Cu84 alloy solidified in drop tube. The size of the droplet is 150 m. The Co-rich L1-phase solidified as sphere and nearly dispersed homogeneously in the entire droplets. The largest diameter of the L1-phase is about 3.3 m. Comparing the samples solidified in EML and DT processing, the electromagnetic stirring and density difference induced stokes motion influences the coalescence of the L1-phase in EML procedure. While in drop tube experiment, both effects are suppressed, the minor phase distributes homogeneously in the matrix phase. 4. Conclusion Both Co41.8 Cu58.2 and Co16 Cu84 alloys were treated by electromagnetic levitation and drop tube processing. Deep
This work is financially supported by the National Natural Science Foundation of China (Grant No. 50221101). One of the authors (X.Y. Lu) is grateful to BMBF for a visiting scholarship. The authors gratefully acknowledge Prof. L. Ratke, Ms. X.R. Liu, Dr. J.R. Gao, Dr. T. Volkmann and Dr. O. Funke for useful discussions and their help with experiments. References [1] Y. Nakagawa, Acta Metall. 6 (1958) 704. [2] S.P. Elder, A. Munitz, G.J. Abbaschian, Mater. Sci. Forum 50 (1989) 137. [3] A. Munitz, R. Abbaschian, Metall. Mater. Trans. 27A (1996) 4049. [4] A. Munitz, S.P. Elder, R. Abbaschian, Metall. Mater. Trans. 23A (1992) 1817. [5] A. Munitz, R. Abbaschian, J. Mater. Sci. 33 (1998) 3639. [6] I. Yamauchi, N. Ueno, M. Shimaoka, I. Ohnaka, J. Mater. Sci. 33 (1998) 371. [7] Z. Sun, X. Song, Z. Hu, S. Yang, G. Liang, J. Sun, J. Alloys Comp. 319 (2001) 266. [8] F. Gillessen, Forschungsbericht DFVLR, DLR-FB 89-32, Köln, 1989. [9] D.M. Herlach, R. Willnecker, F. Gillessen, in: Proceedings of the 5th European Symposium on Materials Science under Microgravity, ESA SP-2 22, European Space Agency, Noordwijk, The Netherlands, 1985, p. 399. [10] T. Nishizawa, K. Ishida, Bull. Alloy Phase Diagrams 5 (1984) 161. [11] C.D. Cao, G.P. Görler, D.M. Herlach, B. Wei, Mater. Sci. Eng. A 325 (2002) 503. [12] C.D. Cao, T. Letzig, G.P. Görler, D.M. Herlach, J. Alloys Comp. 325 (2001) 113. [13] J.R. Rogers, R.H. Davis, Metall. Trans. 21A (1990) 59. [14] L. Ratke, S. Diefenbach, Mater. Sci. Eng. R 15 (1995) 263. [15] R. Eichel, I. Egry, Z. Metallkd. 90 (1999) 371.