- Email: [email protected]
Direct determination of the metastable liquid miscibility gap in undercooled Cu–Co alloys
July 1998
Materials Letters 36 Ž1998. 152–156
Direct determination of the metastable liquid miscibility gap in undercooled Cu–Co alloys D. Li a
a,)
, M.B. Robinson a , T.J. Rathz b, G. Williams
b
Space Sciences Laboratory, NASAr Marshall Space Flight Center, HuntsÕille, AL 35812, USA b UniÕersity of Alabama, HuntsÕille, AL 35899, USA Received 9 January 1998; accepted 13 January 1998
Abstract Bulk Cu–Co alloys at compositions ranging from 10 to 80 wt.% Co were highly undercooled using a melt fluxing technique. The metastable liquid separation boundary has been directly determined from the measured temperature-time profiles. It was also calculated using a subregular solution model. The critical point of the miscibility gap was found to be about 80 K below the liquidus. A droplet-shaped microstructure was observed for all solidified specimens ŽCu–10 to 80 wt.% Co., when the melts were undercooled into the metastable miscibility boundary. q 1998 Elsevier Science B.V. All rights reserved. PACS: 64.70.Ja; 64.75.q g; 81.30.Bx; 81.30.Fb Keywords: Undercooled melts; Liquid separation; Miscibility gap; Cu–Co alloys; Dispersion microstructure; Coarsening
Some alloys exhibit positive deviations from an ideal solution, but they are not large enough to produce a monotectic reaction under equilibrium conditions. However, the deviations result in a flattened liquidus on the phase diagram which implies a thermodynamic tendency to immiscibility upon undercooling DT. Several examples of the submerged liquid miscibility gaps in metallic alloys have been reported w1,2x. A number of Cu-based alloys such as Cu–Co w3x, Cu–Fe w4x, Cu–Nb w5x and Cu–Ta w6x, have long been of interest for the appearance of a miscibility gap in the undercooled liquid. For the former two systems, liquid immiscibility was noticed four decades ago w7x. However, the early paper is limited to low DT near the critical point. Owing to a serious oxidation problem, the possibility cannot be precluded that the separation was induced by impurities. Recently, the microstructures of Cu–10 to 65 wt.% Co samples processed through a combination of containerless undercooling and rapid quenching have been studied w8,9x. Although it is possible to set up a miscibility gap based on both post-mortem composition analysis and sample release temperatures, this method suffers from apparent uncertainties in deciding undercooling and concentration. Moreover, it is difficult to understand that not all samples w9x could form with droplet-shaped microstructure, though they had deep
)
Corresponding author. E-mail: [email protected].
00167-577Xr98r$19.00 q 1998 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 5 7 7 X Ž 9 8 . 0 0 0 3 9 - 1
D. Li et al.r Materials Letters 36 (1998) 152–156
153
excursions into the metastable region. Also the liquid separation was observed in the Cu–Co ribbons w10x produced directly by quenching. As such, it seems ambiguous whether the resulting microstructures obtained from the combined processing w9x are due to the effect of undercooling or cooling rate. To our knowledge, there has been little or no work on direct measurements of submerged liquid miscibility gap in metallic systems, which defies correlation of the metastable demixing with undercooling. Techniques to achieve high undercoolings fall into two categories: containerless processing w11x and melt fluxing w12x. In the former, the levitation force, especially in an electromagnetic field, usually gives rise to a rotation of molten drops and makes it extremely difficult to diagnose melt separation having a small enthalpy. Recent work w13x has demonstrated that comparably large undercoolings have been achieved with melt fluxing, as long as an appropriate flux agent can be found to remove heterogeneous nucleants. As the samples are stationary, this technique allows one to leisurely study the undercooling behavior during slow cooling. In this letter, a substantial degree of undercooling up to 330 K has been achieved for Cu–Co alloys in a flux. For the first time, the metastable miscibility gap has been directly measured in this system with a wide range of compositions. According to a certain ratio, five nine pure copper shots and cobalt pieces Žtotal mass 1.3 g., together with some flux Ža Duran glass. were inserted into an Al 2 O 3 crucible. The crucible was placed in an induction coil connected to a RF generator. In an ultrahigh vacuum ŽpressureF 1.3 = 10y5 Pa. environment, the sample was heated to about 1300 K for 15 min to dehydrate the flux and evaporate other surface impurities. After the in vacuo treatment, the chamber was back-filled with He–6%H 2 gas Žpurity better than 5 N. to 78 kPa. Subsequently, the sample was subjected to a number of successive heating–cooling cycles. The temperature of the open surface of the sample was measured using a single color pyrometer ŽMikron, model 190 with resolution of 0.1 K.. Temperatures were calibrated with the known peritectic transformation. The bulk compositions of specimens were measured by electron microprobe quantitative analysis. The microprobe analysis shows that little has changed in the alloy compositions after processing. The maximum undercooling achievable DT M in the Cu–Co system was found to increase with the Co concentration. The DTM varies from 190 K for Cu–10 wt.% Co to 330 K for Cu–80 wt.% Co alloy. As an example, the cooling profiles of Cu–70 wt.% Co samples with different undercoolings are illustrated in Fig. 1. All the curves manifest relatively low cooling rates of 50 to 70 Krs before solid nucleation and contain two solidification events: growth of primary a-Co dendrites and the peritectic ´-Cu phase. But they differ in the undercooling behavior. The liquidus temperature, TL , can be determined from the first thermal arrest on the curve Ža. of a very small DT. The measured T L coincides with the value on the phase diagram w6x, which verifies the present temperature calibration. For the curve Žb., the melt was smoothly undercooled 130 K below the TL and then led
Fig. 1. Typical cooling curves of Cu–70 wt.% Co samples in flux: Ža. little DT ; Žb. DT s130 K; and Žc. DT s 228 K, the arrow at TS is indicative of liquid separation.
154
D. Li et al.r Materials Letters 36 (1998) 152–156
to a recalescence reflecting the growth of a-Co phase. Further cooled to TP Ž1385 K., the peritectic solidification took place. At still higher undercooling, additional anomalies appeared prior to the major recalescence, as marked in Fig. 1c. This unusual phenomenon neither emerged at small undercooling nor arose from the flux agent of which the glass transition lies below 800 K. Therefore, this precursor reaction is attributed to a phase separation in the highly undercooled alloy. The first angular point prior to recalescence onset is referred to hereafter as TS , the temperature at which separation into two metastable liquids begins to occur. This temperature is reproducible for the specimens of the same composition. Plotting the measured TS on the Cu–Co phase diagram w6x, the metastable liquid miscibility boundary can be established, as presented in Fig. 2. It is of interest to calculate the metastable gap and then compare with the experimental data. Using a subregular solution model, the liquid Gibbs energies of mixing can be expressed as: DG L s V Ž L . x A x B q RT Ž x A ln x A q x B q ln x B .
Ž 1.
where V Ž L., R, x A and x B denote the interaction parameter, gas constant, atom fraction of Cu and Co, respectively. Thermodynamic evaluations of the equilibrium Cu–Co phase diagram have been made by Hasebe and Nishizawa w14x. They put forward a high positive value of interaction parameter in the system: V Ž L. s 1.75 RTmB ŽTmB is the melting point of Co. and V Ž S . s 2.0 RTmB . That is to say, the interaction parameter accrues as the temperature goes down. For simplicity, V Ž L. was treated as a function of temperature T in the range of TP Ž1385. to 1700 K, and a constant beyond this T range in our calculation. Thus, V Ž L. can be modified as:
V Ž L . s 2 y 1.05 = 10y3 = Ž T y Tp . = RTmB
Ž 2.
The DG L was calculated according to Eq. Ž1.. Accordingly, the calculated metastable binodal was obtained through drawing a common tangent on the two negative humps of energy curves. With the exception of a few points, for example, for Cu–40 wt.% Co, the modeling Žthin solid line in Fig. 2. agrees with the measurements Žempty circles.. As compared with the miscibility gap w9x derived from composition analysis on the quenched samples, our directly-determined boundary is depressed about 40 K. In that technique w9x, the drop release temperature was reckoned as the actual solidification temperature, thus overestimating the latter. The metastable liquid separation is further confirmed by the solidified microstructures. By investigation on a whole series of Cu–Co samples which entered the metastable region before solidifying, the droplet-shaped
Fig. 2. The phase diagram of Cu–Co w6x from the liquid to 1300 K. Superposed on it are: the present measurements of metastable melt separation Žempty circles. and calculated miscibility gap Žthin solid line..
D. Li et al.r Materials Letters 36 (1998) 152–156
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microstructures were developed for all compositions from 10 to 80 wt.% Co. Nevertheless, there are two types of particle size distributions: Ža. even dispersions for the Cu or Co-rich alloys and Žb. a steep jump from small droplets to an eminently large one up to 5 mm in diameter for the alloys between 30 and 60 wt.% Co. The promotion of one large droplet in those specimens can be interpreted by the coarsening effect in the final stage of decomposition. It was revealed that the dominant mechanism for causing coarsening under similar processing conditions is related to the volume fraction of the minority component w15x. In the present study, a dispersion can be preserved for the alloys below 30 or above 60 wt.% Co. The microstructure of a Cu–15 wt.% Co specimen undercooled 255 K is shown in Fig. 3a, containing finely dispersed Co-rich spheres with diameters less than 10 m m in the Cu matrix on the entire sample section. Fig. 3b and c display different morphologies for two Cu–70 wt.% Co samples corresponding to cooling curves Žb. and Žc. of Fig. 1, respectively. Alloys solidifying from the homogeneous liquid of small DT possess a ‘normal’ dendritic morphology, as seen in Fig. 3b. In contrast, liquid phase separation at high DT brought about a microstructure with well-defined Cu-rich droplets in the Co-rich matrix shown in Fig. 3c. In this case, the microstructural selection is dictated by the presence of metastable miscibility gap. Another factor influencing the final structure is a possible remixing owing to recalescence. Its maximum temperature TR was found to be dependent on the content of Co. The droplet-shaped structure can be completely retained in the Cu-rich alloys, for the TR is lower than the miscibility gap TS ; while the TR is beyond the TS for the Co-rich alloys, the system is possibly in a miscible state again. However, it is noted that the temperature descended from the TR to TS within two seconds estimated from Fig. 1. This suggests that a massive remixing may be kinetically difficult during this short period, thus keeping the droplet-shaped microstructures. In conclusion, upon solidification of the phase separated liquid in highly undercooled state, droplet-shaped morphologies were formed over a wide interval of compositions from 10 to 80 wt.% Co. Furthermore, uniformly dispersed microstructures were obtained for highly undercooled Cu or Co-rich alloys. The metastable
Fig. 3. Microstructures of: Ža. a Cu–15 wt.% Co sample undercooled 255 K; Žb. and Žc. two Cu–70 wt.% Co samples undercooled 130 and 228 K corresponding to Fig. 1b and c, respectively.
D. Li et al.r Materials Letters 36 (1998) 152–156
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miscibility gap has been directly constructed for the first time in this system. There is good agreement between the measured and calculated metastable liquid miscibility gap.
Acknowledgements This work was performed while one of the authors ŽDL. held a National Research Council-ŽNASA MSFC. Research Associateship. The authors also thank G. Jerman and H. Alexander for composition analysis.
References w1x w2x w3x w4x w5x w6x w7x w8x w9x w10x w11x w12x w13x w14x w15x
J.W. Elmer, M.J. Aziz, L.E. Tanner, P.M. Smith, M.A. Wall, Acta Metall. Mater. 42 Ž1994. 1065. J.H. Perepezko, Mater. Sci. Eng. A 226–228 Ž1997. 374. T. Nishizawa, K. Ishida, Bull. Alloy Phase Diagrams 5 Ž1984. 161. O. Drbohlav, W.J. Filho, A.R. Yavari, Mater. Sci. Forum 225–227 Ž1996. 359. J.D. Verhoeven, H.L. Downing, L.S. Chumbley, E.D. Gibson, J. Appl. Phys. 65 Ž1989. 1293. ASM International, ASM’s Binary Alloy Phase Diagrams on CD-ROM, 2nd edn., 1996. Y. Nakagawa, Acta Metall. 6 Ž1958. 704. S.P. Elder, A. Munitz, G.J. Abbaschian, Mater. Sci. Forum 50 Ž1989. 137. A. Munitz, R. Abbaschian, Metall. Mater. Trans. 27A Ž1996. 4049. J. Wecker, R. Helmolt, L. Schultz, K. Samwer, Appl. Phys. Lett. 62 Ž1993. 1985. R.J. Bayuzick, W.H. Hofmeister, C.W. Morton, M.B. Robinson, in: F. Szofran, D. McCauley, C. Walker ŽEds.., NASA Microgravity Materials Science Conference, NASA Conference Publication 3342, 1996, p. 125. G. Devaud, D. Turnbull, Appl. Phys. Lett. 46 Ž1985. 844. G. Wilde, G.P. Gorler, R. Willnecker, Appl. Phys. Lett. 69 Ž1996. 2995. M. Hasebe, T. Nishizawa, Calphad 4 Ž1980. 83. H.U. Walter, in: E. Rolfe, B. Battrick ŽEds.., The Effect of Gravity on the Solidification of Immiscible Alloys, ESA SP-219, Netherlands, 1984, p. 47.
Materials Letters 36 Ž1998. 152–156
Direct determination of the metastable liquid miscibility gap in undercooled Cu–Co alloys D. Li a
a,)
, M.B. Robinson a , T.J. Rathz b, G. Williams
b
Space Sciences Laboratory, NASAr Marshall Space Flight Center, HuntsÕille, AL 35812, USA b UniÕersity of Alabama, HuntsÕille, AL 35899, USA Received 9 January 1998; accepted 13 January 1998
Abstract Bulk Cu–Co alloys at compositions ranging from 10 to 80 wt.% Co were highly undercooled using a melt fluxing technique. The metastable liquid separation boundary has been directly determined from the measured temperature-time profiles. It was also calculated using a subregular solution model. The critical point of the miscibility gap was found to be about 80 K below the liquidus. A droplet-shaped microstructure was observed for all solidified specimens ŽCu–10 to 80 wt.% Co., when the melts were undercooled into the metastable miscibility boundary. q 1998 Elsevier Science B.V. All rights reserved. PACS: 64.70.Ja; 64.75.q g; 81.30.Bx; 81.30.Fb Keywords: Undercooled melts; Liquid separation; Miscibility gap; Cu–Co alloys; Dispersion microstructure; Coarsening
Some alloys exhibit positive deviations from an ideal solution, but they are not large enough to produce a monotectic reaction under equilibrium conditions. However, the deviations result in a flattened liquidus on the phase diagram which implies a thermodynamic tendency to immiscibility upon undercooling DT. Several examples of the submerged liquid miscibility gaps in metallic alloys have been reported w1,2x. A number of Cu-based alloys such as Cu–Co w3x, Cu–Fe w4x, Cu–Nb w5x and Cu–Ta w6x, have long been of interest for the appearance of a miscibility gap in the undercooled liquid. For the former two systems, liquid immiscibility was noticed four decades ago w7x. However, the early paper is limited to low DT near the critical point. Owing to a serious oxidation problem, the possibility cannot be precluded that the separation was induced by impurities. Recently, the microstructures of Cu–10 to 65 wt.% Co samples processed through a combination of containerless undercooling and rapid quenching have been studied w8,9x. Although it is possible to set up a miscibility gap based on both post-mortem composition analysis and sample release temperatures, this method suffers from apparent uncertainties in deciding undercooling and concentration. Moreover, it is difficult to understand that not all samples w9x could form with droplet-shaped microstructure, though they had deep
)
Corresponding author. E-mail: [email protected].
00167-577Xr98r$19.00 q 1998 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 5 7 7 X Ž 9 8 . 0 0 0 3 9 - 1
D. Li et al.r Materials Letters 36 (1998) 152–156
153
excursions into the metastable region. Also the liquid separation was observed in the Cu–Co ribbons w10x produced directly by quenching. As such, it seems ambiguous whether the resulting microstructures obtained from the combined processing w9x are due to the effect of undercooling or cooling rate. To our knowledge, there has been little or no work on direct measurements of submerged liquid miscibility gap in metallic systems, which defies correlation of the metastable demixing with undercooling. Techniques to achieve high undercoolings fall into two categories: containerless processing w11x and melt fluxing w12x. In the former, the levitation force, especially in an electromagnetic field, usually gives rise to a rotation of molten drops and makes it extremely difficult to diagnose melt separation having a small enthalpy. Recent work w13x has demonstrated that comparably large undercoolings have been achieved with melt fluxing, as long as an appropriate flux agent can be found to remove heterogeneous nucleants. As the samples are stationary, this technique allows one to leisurely study the undercooling behavior during slow cooling. In this letter, a substantial degree of undercooling up to 330 K has been achieved for Cu–Co alloys in a flux. For the first time, the metastable miscibility gap has been directly measured in this system with a wide range of compositions. According to a certain ratio, five nine pure copper shots and cobalt pieces Žtotal mass 1.3 g., together with some flux Ža Duran glass. were inserted into an Al 2 O 3 crucible. The crucible was placed in an induction coil connected to a RF generator. In an ultrahigh vacuum ŽpressureF 1.3 = 10y5 Pa. environment, the sample was heated to about 1300 K for 15 min to dehydrate the flux and evaporate other surface impurities. After the in vacuo treatment, the chamber was back-filled with He–6%H 2 gas Žpurity better than 5 N. to 78 kPa. Subsequently, the sample was subjected to a number of successive heating–cooling cycles. The temperature of the open surface of the sample was measured using a single color pyrometer ŽMikron, model 190 with resolution of 0.1 K.. Temperatures were calibrated with the known peritectic transformation. The bulk compositions of specimens were measured by electron microprobe quantitative analysis. The microprobe analysis shows that little has changed in the alloy compositions after processing. The maximum undercooling achievable DT M in the Cu–Co system was found to increase with the Co concentration. The DTM varies from 190 K for Cu–10 wt.% Co to 330 K for Cu–80 wt.% Co alloy. As an example, the cooling profiles of Cu–70 wt.% Co samples with different undercoolings are illustrated in Fig. 1. All the curves manifest relatively low cooling rates of 50 to 70 Krs before solid nucleation and contain two solidification events: growth of primary a-Co dendrites and the peritectic ´-Cu phase. But they differ in the undercooling behavior. The liquidus temperature, TL , can be determined from the first thermal arrest on the curve Ža. of a very small DT. The measured T L coincides with the value on the phase diagram w6x, which verifies the present temperature calibration. For the curve Žb., the melt was smoothly undercooled 130 K below the TL and then led
Fig. 1. Typical cooling curves of Cu–70 wt.% Co samples in flux: Ža. little DT ; Žb. DT s130 K; and Žc. DT s 228 K, the arrow at TS is indicative of liquid separation.
154
D. Li et al.r Materials Letters 36 (1998) 152–156
to a recalescence reflecting the growth of a-Co phase. Further cooled to TP Ž1385 K., the peritectic solidification took place. At still higher undercooling, additional anomalies appeared prior to the major recalescence, as marked in Fig. 1c. This unusual phenomenon neither emerged at small undercooling nor arose from the flux agent of which the glass transition lies below 800 K. Therefore, this precursor reaction is attributed to a phase separation in the highly undercooled alloy. The first angular point prior to recalescence onset is referred to hereafter as TS , the temperature at which separation into two metastable liquids begins to occur. This temperature is reproducible for the specimens of the same composition. Plotting the measured TS on the Cu–Co phase diagram w6x, the metastable liquid miscibility boundary can be established, as presented in Fig. 2. It is of interest to calculate the metastable gap and then compare with the experimental data. Using a subregular solution model, the liquid Gibbs energies of mixing can be expressed as: DG L s V Ž L . x A x B q RT Ž x A ln x A q x B q ln x B .
Ž 1.
where V Ž L., R, x A and x B denote the interaction parameter, gas constant, atom fraction of Cu and Co, respectively. Thermodynamic evaluations of the equilibrium Cu–Co phase diagram have been made by Hasebe and Nishizawa w14x. They put forward a high positive value of interaction parameter in the system: V Ž L. s 1.75 RTmB ŽTmB is the melting point of Co. and V Ž S . s 2.0 RTmB . That is to say, the interaction parameter accrues as the temperature goes down. For simplicity, V Ž L. was treated as a function of temperature T in the range of TP Ž1385. to 1700 K, and a constant beyond this T range in our calculation. Thus, V Ž L. can be modified as:
V Ž L . s 2 y 1.05 = 10y3 = Ž T y Tp . = RTmB
Ž 2.
The DG L was calculated according to Eq. Ž1.. Accordingly, the calculated metastable binodal was obtained through drawing a common tangent on the two negative humps of energy curves. With the exception of a few points, for example, for Cu–40 wt.% Co, the modeling Žthin solid line in Fig. 2. agrees with the measurements Žempty circles.. As compared with the miscibility gap w9x derived from composition analysis on the quenched samples, our directly-determined boundary is depressed about 40 K. In that technique w9x, the drop release temperature was reckoned as the actual solidification temperature, thus overestimating the latter. The metastable liquid separation is further confirmed by the solidified microstructures. By investigation on a whole series of Cu–Co samples which entered the metastable region before solidifying, the droplet-shaped
Fig. 2. The phase diagram of Cu–Co w6x from the liquid to 1300 K. Superposed on it are: the present measurements of metastable melt separation Žempty circles. and calculated miscibility gap Žthin solid line..
D. Li et al.r Materials Letters 36 (1998) 152–156
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microstructures were developed for all compositions from 10 to 80 wt.% Co. Nevertheless, there are two types of particle size distributions: Ža. even dispersions for the Cu or Co-rich alloys and Žb. a steep jump from small droplets to an eminently large one up to 5 mm in diameter for the alloys between 30 and 60 wt.% Co. The promotion of one large droplet in those specimens can be interpreted by the coarsening effect in the final stage of decomposition. It was revealed that the dominant mechanism for causing coarsening under similar processing conditions is related to the volume fraction of the minority component w15x. In the present study, a dispersion can be preserved for the alloys below 30 or above 60 wt.% Co. The microstructure of a Cu–15 wt.% Co specimen undercooled 255 K is shown in Fig. 3a, containing finely dispersed Co-rich spheres with diameters less than 10 m m in the Cu matrix on the entire sample section. Fig. 3b and c display different morphologies for two Cu–70 wt.% Co samples corresponding to cooling curves Žb. and Žc. of Fig. 1, respectively. Alloys solidifying from the homogeneous liquid of small DT possess a ‘normal’ dendritic morphology, as seen in Fig. 3b. In contrast, liquid phase separation at high DT brought about a microstructure with well-defined Cu-rich droplets in the Co-rich matrix shown in Fig. 3c. In this case, the microstructural selection is dictated by the presence of metastable miscibility gap. Another factor influencing the final structure is a possible remixing owing to recalescence. Its maximum temperature TR was found to be dependent on the content of Co. The droplet-shaped structure can be completely retained in the Cu-rich alloys, for the TR is lower than the miscibility gap TS ; while the TR is beyond the TS for the Co-rich alloys, the system is possibly in a miscible state again. However, it is noted that the temperature descended from the TR to TS within two seconds estimated from Fig. 1. This suggests that a massive remixing may be kinetically difficult during this short period, thus keeping the droplet-shaped microstructures. In conclusion, upon solidification of the phase separated liquid in highly undercooled state, droplet-shaped morphologies were formed over a wide interval of compositions from 10 to 80 wt.% Co. Furthermore, uniformly dispersed microstructures were obtained for highly undercooled Cu or Co-rich alloys. The metastable
Fig. 3. Microstructures of: Ža. a Cu–15 wt.% Co sample undercooled 255 K; Žb. and Žc. two Cu–70 wt.% Co samples undercooled 130 and 228 K corresponding to Fig. 1b and c, respectively.
D. Li et al.r Materials Letters 36 (1998) 152–156
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miscibility gap has been directly constructed for the first time in this system. There is good agreement between the measured and calculated metastable liquid miscibility gap.
Acknowledgements This work was performed while one of the authors ŽDL. held a National Research Council-ŽNASA MSFC. Research Associateship. The authors also thank G. Jerman and H. Alexander for composition analysis.
References w1x w2x w3x w4x w5x w6x w7x w8x w9x w10x w11x w12x w13x w14x w15x
J.W. Elmer, M.J. Aziz, L.E. Tanner, P.M. Smith, M.A. Wall, Acta Metall. Mater. 42 Ž1994. 1065. J.H. Perepezko, Mater. Sci. Eng. A 226–228 Ž1997. 374. T. Nishizawa, K. Ishida, Bull. Alloy Phase Diagrams 5 Ž1984. 161. O. Drbohlav, W.J. Filho, A.R. Yavari, Mater. Sci. Forum 225–227 Ž1996. 359. J.D. Verhoeven, H.L. Downing, L.S. Chumbley, E.D. Gibson, J. Appl. Phys. 65 Ž1989. 1293. ASM International, ASM’s Binary Alloy Phase Diagrams on CD-ROM, 2nd edn., 1996. Y. Nakagawa, Acta Metall. 6 Ž1958. 704. S.P. Elder, A. Munitz, G.J. Abbaschian, Mater. Sci. Forum 50 Ž1989. 137. A. Munitz, R. Abbaschian, Metall. Mater. Trans. 27A Ž1996. 4049. J. Wecker, R. Helmolt, L. Schultz, K. Samwer, Appl. Phys. Lett. 62 Ž1993. 1985. R.J. Bayuzick, W.H. Hofmeister, C.W. Morton, M.B. Robinson, in: F. Szofran, D. McCauley, C. Walker ŽEds.., NASA Microgravity Materials Science Conference, NASA Conference Publication 3342, 1996, p. 125. G. Devaud, D. Turnbull, Appl. Phys. Lett. 46 Ž1985. 844. G. Wilde, G.P. Gorler, R. Willnecker, Appl. Phys. Lett. 69 Ž1996. 2995. M. Hasebe, T. Nishizawa, Calphad 4 Ž1980. 83. H.U. Walter, in: E. Rolfe, B. Battrick ŽEds.., The Effect of Gravity on the Solidification of Immiscible Alloys, ESA SP-219, Netherlands, 1984, p. 47.