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Melt growth processing of RE123 superconductive oxides
PII:
Applied Superconductivity Vol. 4, Nos 10±11, pp. 519±533, 1996 # 1998 Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain 0964-1807/96 $15.00 + 0.00 S0964-1807(97)00037-9
MELT GROWTH PROCESSING OF RE123 SUPERCONDUCTIVE OXIDES K. KAKIMOTO* and Y. SHIOHARA Superconductivity Research Laboratory (SRL), International Superconductivity Technology Center (ISTEC), 1-10-13 Shinonome, Koto-ku, Tokyo 135, Japan AbstractÐThe enhancement of the superconducting properties of the REBa2Cu3O7 ÿ d (RE, rare earth) synthesized by the various methods is reviewed from a crystal growth viewpoint. The diculty in making single crystals is discussed using the phase diagram of the REBa2Cu3O7 ÿ d. In this article, the problem of enhancing the critical current density (Jc) for application to conductors is raised and discussed. In particular, we describe the recent progress of crystal growth research on the RE1 + xBa2 ÿ xCu3O7 ÿ d. # 1998 Elsevier Science Ltd. All rights reserved
NOMENCLATURE C Ce DL Fd Fi G Jc R R Ra Rab Rc RE Rmax R* r* r*a r*c t Tc DT a G Z s Ds0 DsSP DsLP DsSL
practical concentration equilibrium concentration of the solute in the solution diusivity of Y in the liquid drag force the force due to the interfacial energy temperature gradient critical current density growth rate (Section 4) mean radius of Y211 rods (Section 5) growth rate in h100i-direction growth rate in h110i-direction growth rate in h001i-direction rare earth maximum growth rate critical growth rate critical size of a particle critical radius of a Y211 particle for a-direction growth critical radius of a Y211 particle for c-direction growth holding time critical temperature undercooling constant Gibbs±Thomson coecient melt viscosity supersaturation interfacial energy solid±particle interfacial energy liquid±particle interfacial energy solid±liquid interfacial energy
1. INTRODUCTION
Since the discovery of high critical temperature (Tc) superconducting (HTSC) materials by Bednorz and MuÈller [1], tremendous eorts have been made to discover the explore materials with a higher Tc. At last, the Tc of HTSC exceeded the liquid nitrogen temperature with the discovery of YBa2Cu3O7 ÿ d (Y123) materials in 1986. This event stimulated thousands of researchers world-wide, and the researchers studied the possibility of HTSC applications in numerous ®elds from a commercial stand point, since liquid nitrogen, which is much less expensive, easier to handle and has a higher heat capacity than liquid helium, can be used as a coolant. For such *Author to whom correspondence should be addressed. 519
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Fig. 1. The Jc±B curve of Nd123 single crystals measured by the magnetization [2].
application, the high critical current density (Jc) wire and tape materials for the transportation of large currents, the bulk materials for levitation and shielding, thin ®lms for superconducting devices, and so on are given. The ®rst and the second categories are expected immediately in these applications. To realize the ®rst and the second categories, large Jc of the order of 104±106 A cmÿ2 and high Tc values are required in the operating magnetic ®eld. Unfortunately, Jc enhancement was not realized immediately after the discovery of high-temperature superconductivity. This was because Jc is strongly microstructure dependent rather than an intrinsic property of the superconductor. Furthermore, some of the characteristic features of HTSC materials, such as the two-dimensional nature in crystal structure and the very short coherence length, caused severe diculties in Jc enhancement. Hence, the control of the microstructure is very important. In principle, high-Jc applications require a structurally perfect matrix of bulk superconductors, with homogeneously dispersed defects or inclusions acting as magnetic ¯ux pinning sites. Of the many HTSC materials discovered so far, Y123 has been the most widely studied. Various eorts have been made to clarify the formation of the Y123 phase crystals and to improve the Jc property. The mechanism for making Y123 single crystals has gradually become clear. Recently, Nd1 + xBa2 ÿ xCu3O6 + d (Nd123) material drew the attention of the researchers. Figure 1 shows the Jc±B property of Nd123 single crystals measured by the magnetization [2]. As can be seen in this ®gure, Nd123 materials have a signi®cant peak eect in the high magnetic ®eld. For applications in the high magnetic ®eld, Nd123 materials are very attractive. It is, however, dicult to synthesize just Nd123 due to its nonstoichiometric nature. In this review, we describe the recent progress in enhancing the superconducting properties from the viewpoint of the crystal growth of Y123 and other RE123 (RE, rare earth) superconducting oxides.
Fig. 2. Calculated isothermal sections of the Y2O3±BaO±CuOx system at 1223 K (a) and 1273 K (b) and 0.21 atmospheres oxygen pressure [3].
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Fig. 3. Calculated vertical section of the Y2O3±BaO±CuOx system at 0.21 atmospheres oxygen pressure [3].
2. PHASE DIAGRAM OF HTSC MATERIALS
Phase diagrams provide much information to understand the stability of materials under certain conditions. Therefore, phase diagrams are essential as the map to make complicated materials like the RE123. Figure 2(a and b) shows the calculated isothermal sections of the Y2O3± BaO±CuO ternary system at 1223 and 1273 K, respectively, at 0.21 atmospheres oxygen partial pressure reported by Lee and Lee [3, 4]. They performed thermodynamic calculations by ®tting the experimental results of the Y±Ba±Cu±O system of the basis of the results obtained by Lay and Renlund [5] and Lindermer et al. [6]. A comparison of these results indicates that the composition region of two-phase equilibrium between liquid and YBa2Cu3O7 ÿ d is reduced rapidly with increasing temperature because of the evolution of Y2BaCuO5 even in the low Y2O3 concentration region. To show this more eectively, a vertical section containing Y2BaCuO5 (Y211), YBa2Cu3O7 ÿ d (Y123) and a mixture of 3BaCuO2 and 2CuO is shown in Fig. 3. This ®gure shows that the liquidus slope of Y123 is very steep over a small temperature range from 1237 to 1275 K. Furthermore, Y123 and Y211 phases are formed by the peritectic reactions between liquid Y211 and liquid Y2O3, respectively. From the Y±Ba±Cu±O phase diagram, the following characteristic features can be identi®ed. 1. The Y123 tetragonal phase can be formed by a peritectic reaction from solid Y211 and Y-de®cient liquid BaO±CuO. Therefore, it is impossible to grow Y123 crystals by a congruent melt crystal growth procedure, e.g. a simple conventional Czochralski method. 2. The liquidus slope near the Y123 peritectic reaction is very high, which suggests that it is dif®cult to obtain a larger growth rate even if large undercoolings can be applied to the system. 3. The Y concentration in the BaO±CuO liquid is very low. Therefore, the self-¯ux method, which is one of the commonly used processes to produce bulk single crystals before incongruent melting systems, cannot be applied to the production of large single crystals. Most of the rate earth (RE) elements can be substituted for yttrium in the 123 structure. These RE123 phases are formed by peritectic reaction between high-temperature stable RE211 phases and liquid, as in the Y±Ba±Cu±O system; the high-temperature stable phases are, exceptionally, PrBaO4 (Pr110) and Nd4Ba2Cu2O10 (Nd422) in the Pr±Ba±Cu±O and Nd±Ba±Cu±O systems, respectively. The signi®cant feature in the RE±Ba±Cu±O systems is the existence of a solid solution of RE1 + xBa2 ÿ xCu3O7 ÿ d, including light RE elements such as RE 0 La, Pr, Nd, Sm, Eu and Gd, while Y123 is a stoichiometric compound [7±9]. This is considered to be due to the relatively large radii of the RE ions. Figure 4 taken from Kambara et al. [10] shows the Nd123
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Fig. 4. Quasi-ternary phase diagram of the NdO1.5±BaO±CuOx system in air at 10738C near CuOx corner [10].
quasi-ternary phase diagram around the region at 10738C where Nd123 and liquid coexist. It is clear that neither the superconducting Nd123 phase nor the high-temperature stable Nd422 phase is a stoichiometric compound, but solid solutions of Nd1 + xBa2 ÿ xCu3O7 ÿ d and Nd4 + yBa2 ÿ yCu2O10 ÿ d. 3. FUNDAMENTALS OF CRYSTAL GROWTH FROM THE MELT IN HTSC MATERIALS
Figure 5 summarizes the various solidi®cation methods for preparing bulk HTSC crystals. Considering the application and the stability of the melt, these methods may be classi®ed into following two processes: (1) for high-Jc applications, bulk crystals are grown from semi-solids (e.g. RE211 + L); and (2) solidi®cation from the self-¯ux solution is performed for high-crystallinity bulk single crystals. In order to control the microstructures of HTSC materials, we must clarify various phenomena taking place at each stage (liquid, liquid±solid phase transition, solid). Figure 6 and 7 show the dominant phenomena considered to aect the microstructures of HTSC materials in two dierent types of solidi®cation processes, i.e. from semi-solid (i.e. liquid + Y211 phase) and solution (i.e. self-¯ux liquid). Figure 6 illustrates a directional solidi®cation process, such as a zone
Fig. 5. Solidi®cation methods for producing bulk HTSC crystals.
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Fig. 6. Dominant phenomena aecting the microstructure of the Y±Ba±Cu±O system in melt processing.
melting or a Bridgeman method, in the Y±Ba±Cu±O system. Microscopically, the same phenomena take place in nondirectional solidi®cation processing, including melt texture growth (MTG) [11], QMG [12] and MPMG [13] methods. Figure 7 illustrates the SRL-CP method [14], which has developed for pulling a Y123 single crystal and is in fact a modi®ed top-seeded solution growth (TSSG) technique. As shown above, there are many methods to make single crystals. A common conception exists in these methods. In the Y±Ba±Cu±O system, the Y123 crystal grows from the melt. However, it is treated as crystal growth from solution because Y123 melts incongruently and the solubility of yttrium in Ba±Cu±O is very small. In this case, the driving force for crystallization is represented by the supersaturation (s) de®ned as: s C ÿ Ce =Ce ;
1
where C and Ce are the liquid concentration and the equilibrium concentration of the solute in the solution. The crystal growth may be limited by the kinetics of atom attachment to the interface, diusion of solute from the environmental liquid phase and/or diusion of latent heat for crystallization.
Fig. 7. Dominant phenomena aecting the size and crystallinity of Y123 single crystal in the TSSG method (e.g. SRL-CP method).
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Fig. 8. Eect of the temperature gradient G and growth rate R on the morphology of the growing Y123±liquid interface, showing that a higher G/R ratio (greater than 33008C h cmÿ2) is required to obtain planar interfaces. The temperature gradients at the interface used for the samples with Ag2O addition are calculated values.
4. CONTROL OF THE GROWTH DIRECTION
For applications to conductors, the morphology of the solidi®cation interface is a very important factor in the crystal growth from semi-solid because the ideal conductor requires no grain boundary and a highly oriented texture. In general, the morphology of the solidi®cation
Fig. 9. G/R dependence of the Jc value [18].
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Fig. 10. The relationship between the orientation along the ab-planes and the pulling rate evaluated by pole ®gure analysis for the samples with dierent nominal compositions [19].
interface is aected by the temperature gradient (G) and growth rate (R). Figure 8 summarizes the eect of G and R on the morphology of the solidi®cation front, including the results of zone melting and Bridgeman methods investigated by three groups [15±17]. Figure 9, taken from Izumi et al. [18], shows the G/R dependence of the Jc values at the same cooling rate (G R). As this result, an increase in Jc values of zone-melted samples with increasing G/R can be recognized even at the same cooling rates (G R) because of the aligned grain boundaries. In general, the transport current along the direction of the ab-plane is much larger than that along the direction of the c-axis. Therefore, the control of growth direction for the direction of electric current is very important in directional solidi®cation. Figure 10, taken from Imagawa et al. [19], shows the pulling rate dependence of the angle between the growth direction and the orientation along the ab-plane. It is evident that generally a higher pulling rate leads to a better orientation of the ab-plane along the growth direction. Figure 11 shows the schematic illustration of the longitudinal interface in the unidirectional solidi®cation of Y123 materials. The maximum growth rate, Rmax, equals the pulling rate making the rod sample in a steady state growth. It is clearly understood from this ®gure that Ra, Rc and the ratio Ra/Rc could be calculated by using a set of the following equations for each grown sample: p p 2 Ra Rab = 2 Rmax cos y= 2; p Rc Rmax sin y= 2;
3
p Ra =Rc cot y= 2;
4
where Ra, Rab, Rc are the growth rates in h100i-, h110i- and h001i-directions, respectively. Therefore, y (the angle between the ab-plane and the growth direction) is determined by the ratio Ra/Rc. Recently, Endo et al. [20] found that Ra and Rc showed an increasing trend with increasing undercooling (D T) and Rc was faster than Ra within their experimental conditions of low undercooling solidi®cation. They also revealed that D T followed the parabolic equation, i.e. RaADT1.9 and RcADT1.3, respectively. These results are shown in Fig. 12 [20]. So it seems that it is possible to vary the orientation of the crystal growth direction and to adjust the ab-plane
Fig. 11. Schematic illustration of the longitudinal interface of a typical sample.
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Fig. 12. Dependence of the growth rate on D T for a-axis and c-axis [20].
orientation close to the crystal growth direction in the undercooling method as well as in directional solidi®cation. In addition, Imagawa et al. [21] succeeded by controlling the growth direction in fabrication of the current lead with the Jc value of 7.0 104 A cmÿ2 (77 K, 0 T) by the conventional four-probe method. 5. DISPERSION OF THE FINE NONSUPERCONDUCTING PARTICLES
For application in high magnetic ®elds, the introduction of the pinning centers into the RE123 materials as the matrix is very important because the samples made by the processes of QMG, MPMG, etc. possess good properties in high magnetic ®elds. Therefore, understanding the mechanism for the introduction of the pinning centers into the RE123 is essential. In this section, size reduction of the particles as the pinning centers and introduction of these pinning centers are described. 5.1. Size reduction of nonsuperconducting particles Understanding of the coarsening behavior of the high-temperature stable phase particles in the liquid is extremely important since ®ner high-temperature stable phase particles are required to enhance magnetic ¯ux pinning as well as RE123 formation through the peritectic reaction. Izumi et al. [22] have investigated the coarsening rate of Y211 phase particles in the liquid by observations of pellet samples held at 10708C for dierent times. Figure 13 shows the holding-
Fig. 13. Holding time dependence of the mean radius of samples with and without Pt doping in a logarithmic plot [22].
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Fig. 14. Histogram of the Y211 distribution [27].
time dependence of the mean radii of Y211 phase particles in a logarithmic coordinate, assuming that the shape of the Y211 phase crystal is spherical, although the Y211 phase forms a rodlike shape crystal with some facets. The gradient of the lines is about the one-third, i.e. the mean radius increases proportionally to t1/3 (t, holding time). According to the above results, the mean radius of the Y211 phase particles increases proportionally to t1/3. Such coarsening behavior (one-third law) is often explained by the Ostwald ripening theory [23]. Izumi et al. [22] calculated the coarsening rate of the rod-like shape by modifying the theory. With several assumptions, the mean radius (R) of Y211 rods can be described as: R a DL Gt1=3 ;
5
where a is a constant, DL is the diusivity of Y in the liquid and G is the Gibbs±Thomson coef®cient. This equation has the same time dependence as that predicted by the Ostwald ripening theory and explains the experimental results of the Y211 coarsening behavior. The change in the coarsening rate is related to the change in the product DLG by Equation (5). Therefore, the eect of Pt doping may be explained by considering the fact that the dissolved solute changes the value of DLG which results in a change in the coarsening rate. By comparing the coarsening rates with and without Pt doping, it was found that the DLG value with platinum addition is about 10% of the value without Pt doping. As mentioned above, the size reduction of nonsuperconductor particles by doping an additive is eective. Recently, as the other examples, Kim et al. [24]
Fig. 15. Schematic drawings showing a particle in front of the solid±liquid interface and the necessary condition of particle pushing: (a) for Ds0>0, the force (Fi) due to the interfacial energy (Ds0) is conducive to pushing; (b) for Ds0<0, Fi is conducive to trapping. The drag force (Fd) is always conducive to trapping.
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Fig. 16. Schematically illustrated pushing/entrapment relationship between the critical growth rate (R*) and the critical radius (r*) of inclusion for dierent growth directions.
and Matsuoka et al. [25] presented that CeO2 is also eective as the additive for the Nd123 materials. As in the other methods, Kambara et al. [26] demonstrated that the size of nonsuperconductive particles is reduced by the Ba-rich compositions in the Nd123 materials. 5.2. Pushing/trapping for HTSC materials In this section, we present how to disperse ®ne high-temperature stable phase inclusions to the matrix of RE123. Figure 14 taken from Endo et al. [27] shows the histogram of the Y211 distribution for the Y123 crystal grown at two dierent undercooling, D T, values. From a quantitative analysis of the Y211 particles in the four regions with dierent D T and dierent growth directions, the number of Y211 particles in the D T = 30 K region is clearly much larger than that in the D T = 10 K region over a whole range of diameters. It should be noted that these phenomena are similar to a pushing/trapping behavior of foreign particles at an advancing solid±liquid interface during solidi®cation. Figure 15 is a schematic drawing showing a particle in front of the solid±liquid interface and the forces acting on the particle. There are two dominant forces: the drag force (Fd) due to viscous ¯ow around the particle, which moves together with the interface at R relative to the melt; and the force (Fi) due to the interfacial energy (Ds0). The other kinds of forces should also strictly be considered, such as gravity [28] and the van der Waals' force [29]. Furthermore, these forces should be dependent on the shape of the solid± liquid interface, which is determined by many factors, such as the thermal conductivity, pressure [30], distribution coecient [31], etc. To simplify the following discussion, however, we assume that the interface shape is planar while treating the two dominant forces. The necessary condition for particle pushing can be described as shown in Equation (6) [32]: Ds0 DsSP ÿ DsLP ÿ DsSL > 0;
6
where DsSP, DsLP and DsSL are the solid±particle, liquid±particle and solid±liquid interfacial energies, respectively. Fi is conductive to pushing if Ds0>0, while Fd is always conducive to trapping as shown in Fig. 15. According to the pushing/trapping theory, the critical size (r*) of a particle, larger than that trapped by the solid, is roughly determined by the critical growth rate (R*) and interfacial energy (Ds0): R ADs0 =Zr
7
where Z is the melt viscosity. Figure 16 shows the schematic relationship between the growth rate and radius of inclusion, which is derived using Equation (6) and considering R* r* = constant. From this ®gure, In the case of D T = 10 K, the critical radius (r*a) of a Y211 particle for the a-direction growth is smaller than r*c because of the anisotropy of R and Ds0. The parameters r*a and r*c at D T = 10 K are considered to be relatively large in the size distribution of Y211 dispersed in front of the interface because the total volume fraction of Y211 particles is much less than the calculated values from the lever rule using the initial composition for both directions. On changing D T from 10 to 30 K, the growth rates for both directions
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Fig. 17. Schematic illustration of pushing/trapping model for the Y123±Y211 system.
increase, leading to a signi®cant decrease in both r*a and r*c . Accordingly, even smaller Y211 particles, which were pushed at D T = 10 K, are trapped by Y123 crystals at D T = 30 K. The above mentioned phenomena are illustrated schematically in Fig. 17. In addition, Endo et al. [33] succeeded by controlling in fabrication of the single crystal bulk with the Jc value of 2 104 A cmÿ2 (77 K, 0.8 T) by magnetization. 6. HOMOGENEITY OF SINGLE CRYSTALS
To enhance the transport current, the larger volume of the RE123 as the matrix should be maintained in the conductor. In the case of the Y123 materials, the balance between the Y123 as the matrix and the Y211 as the pinning center is important because of a stoichiometric compound. However, in the cased of the other RE123 the additional problem of the solid solubility should be studied because of a nonstoichiometric compound. In this section, we describe the recent progress to minimize the solid solubility in the RE123 materials in contrast to the Y123 material. 6.1. Eect of oxygen partial pressure Nakamura et al. [34] reported the eect of oxygen partial pressure on the top-seeded solutiongrowth. The composition of the Nd123 single crystals grown in dierent oxygen partial pressure atmospheres was analyzed by ICP-AES and the results are shown in Table 1. From Table 1, it is recognized that with the same composition of solvent and solute, the deviation from the nominal composition (Nd:Ba:Cu = 1:2:3) is found to increase with increasing oxygen partial pressure atmosphere during crystal growth. This result indicates that a low oxygen partial pressure atmosphere during Nd123 crystal growth is found to be eective for minimizing the substitution of Nd ion into Ba sites. Figure 18 shows the temperature dependence of magnetization for the Nd123 single crystals shown in Table 1. It is obvious that it is possible to enhance the Tc value by minimizing the substitution of Nd ion into Ba sites, and in the research ®elds of the melt processing. Yoo et al. [35] found similar results. Recently, Yoshizumi et al. [36] reported the interesting results as the following. Figure 19 shows NdBCO quasi-ternary phase diagram in low pO2. This phase diagram diers from that in air shown in Fig. 20 [10], but resembles that of YBCO system. The solubility limit of Nd1 + xBa2 ÿ xCu3O7 ÿ d is very small (xmax=0.04). Accordingly, Nd123 superconductors made in the low oxygen partial pressure atmosphere are nearly stoichiometric, and this is the reason why the single crystals grown in the low pO2 atmosphere by Nakamura et al. show high a Tc of 96 K.
Table 1. Composition of Nd123 single crystals grown in dierent oxygen partial pressure atmospheres [34] p(O2) (atm) 0.01 0.21 1.00
Nd:Ba:Cu 1.01:1.97:3.00 1.07:1.95:3.00 1.10:1.90:3.00
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Fig. 18. Temperature dependence of the normalized magnetic susceptibility for Nd123 single crystals grown in dierent oxygen partial pressure atmosphere [34].
6.2. Control of composition Recently, it is reported to minimize the substitution of RE ion into Ba sites by control of the composition. As an example, Fig. 21 by Yao et al. [37] shows the temperature dependence of magnetization of Nd123 single crystals grown from the liquid with dierent ratios of Ba/Cu. The eects of the Ba/Cu ratio in the liquid on Tc and superconducting transition behaviors are clear. With increase of the Ba/Cu ratio, Tc increases and the transition width D T becomes narrow, indicating that the ratio of Ba to Cu in the liquid plays a signi®cant role in controlling the Tc of Nd123 superconductor. From Fig. 4 (see Section 2), it can be seen that points on the liquidus through the tie lines connect to dierent points on the solid solution line. In other words, the substitution of Nd at the Ba site is a function of the Ba/Cu ratio in the liquid. With increase of Ba/Cu ratio in the liquid, the Nd substitution content decreases. When the Ba/Cu is
Fig. 19. NdO1.5±BaO±CuO quasi-ternary equilibrium phase diagram in 1% pO2 at about 1273 K [36].
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Fig. 20. NdO1.5±BaO±CuO quasi-ternary equilibrium phase diagram in air at about 1323 K [10].
approximately 0.8, Nd123 is close to the stoichiometric 123 composition. The Tc results are explained in accordance with this phase diagram. Also in the ®eld of other melt growth, Watanabe et al. [38] and Kambara et al. [26] reported similar results for the bulk materials. 6.3. Eect of growth temperature Kambara et al. [10] investigated the NdBCO quasi-ternary phase diagram for the region of dierent temperatures where Nd123 and liquid coexist in air. According to their results, it was found that the tie lines between Nd123 with the low substitution and the liquid phase lean toward the low substitution as the equilibrium temperature decreases. Therefore, it is considered that the Nd123 with the lower substitution is obtained from the liquid phase of the same composition under the lower temperature. As an example, Fig. 22 shows the relationship between
Fig. 21. Temperature dependence of magnetization for NdBCO single crystals grown in air with dierent ratios of Ba to Cu in the liquid [37].
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Fig. 22. The relationship between the growth temperature and x value in Nd1 + xBa2 ÿ xCu3Oy analyzed by EPMA [39]. The ¯ux composition of Ba/Cu = 3/5 was used.
the growth temperature and the substitution of Nd ion into Ba sites. This result is from Ishida et al. [39], in which the grown ®lm compositions are analyzed for the samples prepared by the LPE method with the ¯ux of Ba/Cu = 3/5. From the result, it is clear to minimize the substitution of Nd ion at the Ba site by preparation under a lower growth temperature, which agrees with the report by Kambara et al. [10]. 7. CONCLUSIONS
The enhancement of superconducting properties in the various methods for the crystal growth of the REBa2Cu3O7 ÿ d has been reviewed focusing on the recent progress in research. The study for the mechanism of the crystal growth of RE123 materials proceeded and the superconducting property also progressed. In the case of application to a conductor under a high magnetic ®eld, the introduction of pinning centers is very important and understanding the crystal growth mechanism becomes more eective. Recently, the Nd123 materials drew the attention of researchers because of its good properties in the high magnetic ®elds and their tremendous eorts have concentrated on the problem for solid solubility in the Nd123 materials. Furthermore, we believe that clari®cation of the growth mechanism of the RE123 materials will assist in further understanding of other oxide materials as well. AcknowledgementsÐThe authors wish to express their sincere gratitude to Dr. A. Endo for his achievement, to Dr. Kambara for making the eective phase diagram, to the other researchers in Division 4 of SRL-ISTEC for kindly providing their results. Parts of the studies presented in this paper were supported by the New Energy and Industrial Technology Development Organization for the R and D Industrial Science and Technology Frontier Program.
REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12.
J. G. Bednorz and K. A. MuÈller, Z. Phys. B 64, 189 (1986). M. Nakamura, Y. Yamada, T. Hirayama, Y. Ikuhara, Y. Shiohara and S. Tanaka, Phys. C 259, 295 (1996). B. J. Lee and D. N. Lee, J. Am. Ceram. Soc. 74, 78 (1991). B. J. Lee and D. N. Lee, J. Am. Ceram. Soc. 72, 314 (1989). K. W. Lay and G. M. Renlund, J. Am. Ceram. Soc. 73, 1208 (1990). T. B. Lindemer, J. F. Hunley, J. E. Gates Jr, A. L. Sutton, J. Brynestad, C. R. Hubbard and P. K. Gallagher, J. Am. Ceram. Soc. 72, 1775 (1989). S. I. Yoo and R. W. McCallum, Phys. C 210, 147 (1993). T. Wada, N. Suzuki, T. Maeda, A. Maeda, S. Uchida, K. Uchinokura and S. Tanaka, Appl. Phys. Lett. 52, 1989 (1988). M. Daeumling, J. M. Seuntjens and D. C. Larbalestier, Nature 346, 332 (1990). M. Kambara, M. Tagami, X. Yao, T. Umeda and Y. Shiohara, J. Am. Ceram. Soc. submitted, (1997). S. Jin, T. H. Tiefel, R. C. Sherwood, R. B. van Dover, M. E. Davis, G. W. Kammlott and R. A. Fastnacht, Phys. Rev. B 37, 7850 (1988). M. Murakami, M. Morita, K. Doi and K. Miyamoto, Jap. J. Appl. Phys. 28, 1189 (1989).
Melt growth processing of superconductive oxides 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39.
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H. Fujimoto, M. Murakami, S. Gotoh, Y. Shiohara, N. Koshizuka and S. Tanaka, Adv. Supercond. 1, 285 (1989). Y. Yamada and Y. Shiohara, Phys. C 217, 182 (1994). T. Izumi, Y. Nakamura and Y. Shiohara, J. Cryst. Growth 128, 757±761 (1993). M. J. Cima, M. C. Flemings, A. M. Figuredo, M. Nakade, H. Ishii, H. D. Brody and J. S. Haggerty, J. Appl. Phys. 72, 179 (1992). N. Mori, H. Hata and K. Ogi, J. Jap. Inst. Met. 6, 648 (1992). T. Izumi and Y. Shiohara, J. Mater. Res. 7, 16 (1992). Y. Imagawa and Y. Shiohara, Phys. C 262, 243 (1996). A. Endo, H. S. Chauhan, Y. Nakamura and Y. Shiohara, J. Mater. Res. 11, 1114 (1996). Y. Imagawa, K. Kakimoto and Y. Shiohara, Phys. C 280, 245 (1997). T. Izumi, Y. Nakamura and Y. Shiohara, J. Mater. Res. 8, (1993). W. D. Kingery, H. K. Bowen and D. R. Uhlmann, Introduction to Ceramics, Wiley, New York, (1960). C. J. Kim, H. W. Park, K. B. Kim, K. W. Lee, I. H. Kuk and G. W. Hong, Mater. Lett. 29, 7 (1996). S. Matsuoka, M. Sumida, T. Umeda and Y. Shiohara, Phys. C 282±287, 531 (1997). M. Kambara, A. Endo, Y. Watanabe, Y. Shiohara and T. Umeda, International Workshop on Superconductivity. p. 202 (1997). A. Endo, H. S. Chauhan, T. Egi and Y. Shiohara, J. Mater. Res. 11, 795 (1996). S. G. Grigoryan, A. S. Oganesyan and A. G. Sarkisyan, Kristallogra®ya 28, 782 (1983). J. Potoschke and V. Rogge, J. Cryst. Growth 94, 726 (1989). G. F. Bolling and J. CisseÂ, J. Cryst. Growth 10, 726 (1971). A. A. Chernov, D. E. Temkin and A. M. Mel'nikova, Sov. Phys. Crystallogr. 21, 369 (1976). D. R. Uhlmann, B. Chalmers and K. A. Jackson, J. Appl. Phys. 35, 2986 (1964). A. Endo, H. S. Chauhan and Y. Shiohara, Phys. C 273, 107 (1996). M. Nakamura, H. Kutami and Y. Shiohara, Phys. C 260, 297 (1996). S. I. Yoo, N. Sakai, H. Takaichi, T. Higuchi and M. Murakami, Appl. Phys. Lett. 65, 633 (1994). M. Yoshizumi, M. Kambara, Y. Shiohara and T. Umeda, International Workshop on Superconductivity. p. 295 (1997). X. Yao, M. Kambara, T. Umeda and Y. Shiohara, Jap. J. Appl. Phys. 36, 400 (1997). Y. Watanabe, K. Miyake, A. Endo, K. Murata, Y. Shiohara and T. Umeda, Phys. C 280, 215 (1997). Y. Ishida, K. Kakimoto, Y. Shiohara and A. B. Sawaoka, Phys. C., (to be submitted).
Applied Superconductivity Vol. 4, Nos 10±11, pp. 519±533, 1996 # 1998 Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain 0964-1807/96 $15.00 + 0.00 S0964-1807(97)00037-9
MELT GROWTH PROCESSING OF RE123 SUPERCONDUCTIVE OXIDES K. KAKIMOTO* and Y. SHIOHARA Superconductivity Research Laboratory (SRL), International Superconductivity Technology Center (ISTEC), 1-10-13 Shinonome, Koto-ku, Tokyo 135, Japan AbstractÐThe enhancement of the superconducting properties of the REBa2Cu3O7 ÿ d (RE, rare earth) synthesized by the various methods is reviewed from a crystal growth viewpoint. The diculty in making single crystals is discussed using the phase diagram of the REBa2Cu3O7 ÿ d. In this article, the problem of enhancing the critical current density (Jc) for application to conductors is raised and discussed. In particular, we describe the recent progress of crystal growth research on the RE1 + xBa2 ÿ xCu3O7 ÿ d. # 1998 Elsevier Science Ltd. All rights reserved
NOMENCLATURE C Ce DL Fd Fi G Jc R R Ra Rab Rc RE Rmax R* r* r*a r*c t Tc DT a G Z s Ds0 DsSP DsLP DsSL
practical concentration equilibrium concentration of the solute in the solution diusivity of Y in the liquid drag force the force due to the interfacial energy temperature gradient critical current density growth rate (Section 4) mean radius of Y211 rods (Section 5) growth rate in h100i-direction growth rate in h110i-direction growth rate in h001i-direction rare earth maximum growth rate critical growth rate critical size of a particle critical radius of a Y211 particle for a-direction growth critical radius of a Y211 particle for c-direction growth holding time critical temperature undercooling constant Gibbs±Thomson coecient melt viscosity supersaturation interfacial energy solid±particle interfacial energy liquid±particle interfacial energy solid±liquid interfacial energy
1. INTRODUCTION
Since the discovery of high critical temperature (Tc) superconducting (HTSC) materials by Bednorz and MuÈller [1], tremendous eorts have been made to discover the explore materials with a higher Tc. At last, the Tc of HTSC exceeded the liquid nitrogen temperature with the discovery of YBa2Cu3O7 ÿ d (Y123) materials in 1986. This event stimulated thousands of researchers world-wide, and the researchers studied the possibility of HTSC applications in numerous ®elds from a commercial stand point, since liquid nitrogen, which is much less expensive, easier to handle and has a higher heat capacity than liquid helium, can be used as a coolant. For such *Author to whom correspondence should be addressed. 519
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Fig. 1. The Jc±B curve of Nd123 single crystals measured by the magnetization [2].
application, the high critical current density (Jc) wire and tape materials for the transportation of large currents, the bulk materials for levitation and shielding, thin ®lms for superconducting devices, and so on are given. The ®rst and the second categories are expected immediately in these applications. To realize the ®rst and the second categories, large Jc of the order of 104±106 A cmÿ2 and high Tc values are required in the operating magnetic ®eld. Unfortunately, Jc enhancement was not realized immediately after the discovery of high-temperature superconductivity. This was because Jc is strongly microstructure dependent rather than an intrinsic property of the superconductor. Furthermore, some of the characteristic features of HTSC materials, such as the two-dimensional nature in crystal structure and the very short coherence length, caused severe diculties in Jc enhancement. Hence, the control of the microstructure is very important. In principle, high-Jc applications require a structurally perfect matrix of bulk superconductors, with homogeneously dispersed defects or inclusions acting as magnetic ¯ux pinning sites. Of the many HTSC materials discovered so far, Y123 has been the most widely studied. Various eorts have been made to clarify the formation of the Y123 phase crystals and to improve the Jc property. The mechanism for making Y123 single crystals has gradually become clear. Recently, Nd1 + xBa2 ÿ xCu3O6 + d (Nd123) material drew the attention of the researchers. Figure 1 shows the Jc±B property of Nd123 single crystals measured by the magnetization [2]. As can be seen in this ®gure, Nd123 materials have a signi®cant peak eect in the high magnetic ®eld. For applications in the high magnetic ®eld, Nd123 materials are very attractive. It is, however, dicult to synthesize just Nd123 due to its nonstoichiometric nature. In this review, we describe the recent progress in enhancing the superconducting properties from the viewpoint of the crystal growth of Y123 and other RE123 (RE, rare earth) superconducting oxides.
Fig. 2. Calculated isothermal sections of the Y2O3±BaO±CuOx system at 1223 K (a) and 1273 K (b) and 0.21 atmospheres oxygen pressure [3].
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Fig. 3. Calculated vertical section of the Y2O3±BaO±CuOx system at 0.21 atmospheres oxygen pressure [3].
2. PHASE DIAGRAM OF HTSC MATERIALS
Phase diagrams provide much information to understand the stability of materials under certain conditions. Therefore, phase diagrams are essential as the map to make complicated materials like the RE123. Figure 2(a and b) shows the calculated isothermal sections of the Y2O3± BaO±CuO ternary system at 1223 and 1273 K, respectively, at 0.21 atmospheres oxygen partial pressure reported by Lee and Lee [3, 4]. They performed thermodynamic calculations by ®tting the experimental results of the Y±Ba±Cu±O system of the basis of the results obtained by Lay and Renlund [5] and Lindermer et al. [6]. A comparison of these results indicates that the composition region of two-phase equilibrium between liquid and YBa2Cu3O7 ÿ d is reduced rapidly with increasing temperature because of the evolution of Y2BaCuO5 even in the low Y2O3 concentration region. To show this more eectively, a vertical section containing Y2BaCuO5 (Y211), YBa2Cu3O7 ÿ d (Y123) and a mixture of 3BaCuO2 and 2CuO is shown in Fig. 3. This ®gure shows that the liquidus slope of Y123 is very steep over a small temperature range from 1237 to 1275 K. Furthermore, Y123 and Y211 phases are formed by the peritectic reactions between liquid Y211 and liquid Y2O3, respectively. From the Y±Ba±Cu±O phase diagram, the following characteristic features can be identi®ed. 1. The Y123 tetragonal phase can be formed by a peritectic reaction from solid Y211 and Y-de®cient liquid BaO±CuO. Therefore, it is impossible to grow Y123 crystals by a congruent melt crystal growth procedure, e.g. a simple conventional Czochralski method. 2. The liquidus slope near the Y123 peritectic reaction is very high, which suggests that it is dif®cult to obtain a larger growth rate even if large undercoolings can be applied to the system. 3. The Y concentration in the BaO±CuO liquid is very low. Therefore, the self-¯ux method, which is one of the commonly used processes to produce bulk single crystals before incongruent melting systems, cannot be applied to the production of large single crystals. Most of the rate earth (RE) elements can be substituted for yttrium in the 123 structure. These RE123 phases are formed by peritectic reaction between high-temperature stable RE211 phases and liquid, as in the Y±Ba±Cu±O system; the high-temperature stable phases are, exceptionally, PrBaO4 (Pr110) and Nd4Ba2Cu2O10 (Nd422) in the Pr±Ba±Cu±O and Nd±Ba±Cu±O systems, respectively. The signi®cant feature in the RE±Ba±Cu±O systems is the existence of a solid solution of RE1 + xBa2 ÿ xCu3O7 ÿ d, including light RE elements such as RE 0 La, Pr, Nd, Sm, Eu and Gd, while Y123 is a stoichiometric compound [7±9]. This is considered to be due to the relatively large radii of the RE ions. Figure 4 taken from Kambara et al. [10] shows the Nd123
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Fig. 4. Quasi-ternary phase diagram of the NdO1.5±BaO±CuOx system in air at 10738C near CuOx corner [10].
quasi-ternary phase diagram around the region at 10738C where Nd123 and liquid coexist. It is clear that neither the superconducting Nd123 phase nor the high-temperature stable Nd422 phase is a stoichiometric compound, but solid solutions of Nd1 + xBa2 ÿ xCu3O7 ÿ d and Nd4 + yBa2 ÿ yCu2O10 ÿ d. 3. FUNDAMENTALS OF CRYSTAL GROWTH FROM THE MELT IN HTSC MATERIALS
Figure 5 summarizes the various solidi®cation methods for preparing bulk HTSC crystals. Considering the application and the stability of the melt, these methods may be classi®ed into following two processes: (1) for high-Jc applications, bulk crystals are grown from semi-solids (e.g. RE211 + L); and (2) solidi®cation from the self-¯ux solution is performed for high-crystallinity bulk single crystals. In order to control the microstructures of HTSC materials, we must clarify various phenomena taking place at each stage (liquid, liquid±solid phase transition, solid). Figure 6 and 7 show the dominant phenomena considered to aect the microstructures of HTSC materials in two dierent types of solidi®cation processes, i.e. from semi-solid (i.e. liquid + Y211 phase) and solution (i.e. self-¯ux liquid). Figure 6 illustrates a directional solidi®cation process, such as a zone
Fig. 5. Solidi®cation methods for producing bulk HTSC crystals.
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Fig. 6. Dominant phenomena aecting the microstructure of the Y±Ba±Cu±O system in melt processing.
melting or a Bridgeman method, in the Y±Ba±Cu±O system. Microscopically, the same phenomena take place in nondirectional solidi®cation processing, including melt texture growth (MTG) [11], QMG [12] and MPMG [13] methods. Figure 7 illustrates the SRL-CP method [14], which has developed for pulling a Y123 single crystal and is in fact a modi®ed top-seeded solution growth (TSSG) technique. As shown above, there are many methods to make single crystals. A common conception exists in these methods. In the Y±Ba±Cu±O system, the Y123 crystal grows from the melt. However, it is treated as crystal growth from solution because Y123 melts incongruently and the solubility of yttrium in Ba±Cu±O is very small. In this case, the driving force for crystallization is represented by the supersaturation (s) de®ned as: s C ÿ Ce =Ce ;
1
where C and Ce are the liquid concentration and the equilibrium concentration of the solute in the solution. The crystal growth may be limited by the kinetics of atom attachment to the interface, diusion of solute from the environmental liquid phase and/or diusion of latent heat for crystallization.
Fig. 7. Dominant phenomena aecting the size and crystallinity of Y123 single crystal in the TSSG method (e.g. SRL-CP method).
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Fig. 8. Eect of the temperature gradient G and growth rate R on the morphology of the growing Y123±liquid interface, showing that a higher G/R ratio (greater than 33008C h cmÿ2) is required to obtain planar interfaces. The temperature gradients at the interface used for the samples with Ag2O addition are calculated values.
4. CONTROL OF THE GROWTH DIRECTION
For applications to conductors, the morphology of the solidi®cation interface is a very important factor in the crystal growth from semi-solid because the ideal conductor requires no grain boundary and a highly oriented texture. In general, the morphology of the solidi®cation
Fig. 9. G/R dependence of the Jc value [18].
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Fig. 10. The relationship between the orientation along the ab-planes and the pulling rate evaluated by pole ®gure analysis for the samples with dierent nominal compositions [19].
interface is aected by the temperature gradient (G) and growth rate (R). Figure 8 summarizes the eect of G and R on the morphology of the solidi®cation front, including the results of zone melting and Bridgeman methods investigated by three groups [15±17]. Figure 9, taken from Izumi et al. [18], shows the G/R dependence of the Jc values at the same cooling rate (G R). As this result, an increase in Jc values of zone-melted samples with increasing G/R can be recognized even at the same cooling rates (G R) because of the aligned grain boundaries. In general, the transport current along the direction of the ab-plane is much larger than that along the direction of the c-axis. Therefore, the control of growth direction for the direction of electric current is very important in directional solidi®cation. Figure 10, taken from Imagawa et al. [19], shows the pulling rate dependence of the angle between the growth direction and the orientation along the ab-plane. It is evident that generally a higher pulling rate leads to a better orientation of the ab-plane along the growth direction. Figure 11 shows the schematic illustration of the longitudinal interface in the unidirectional solidi®cation of Y123 materials. The maximum growth rate, Rmax, equals the pulling rate making the rod sample in a steady state growth. It is clearly understood from this ®gure that Ra, Rc and the ratio Ra/Rc could be calculated by using a set of the following equations for each grown sample: p p 2 Ra Rab = 2 Rmax cos y= 2; p Rc Rmax sin y= 2;
3
p Ra =Rc cot y= 2;
4
where Ra, Rab, Rc are the growth rates in h100i-, h110i- and h001i-directions, respectively. Therefore, y (the angle between the ab-plane and the growth direction) is determined by the ratio Ra/Rc. Recently, Endo et al. [20] found that Ra and Rc showed an increasing trend with increasing undercooling (D T) and Rc was faster than Ra within their experimental conditions of low undercooling solidi®cation. They also revealed that D T followed the parabolic equation, i.e. RaADT1.9 and RcADT1.3, respectively. These results are shown in Fig. 12 [20]. So it seems that it is possible to vary the orientation of the crystal growth direction and to adjust the ab-plane
Fig. 11. Schematic illustration of the longitudinal interface of a typical sample.
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Fig. 12. Dependence of the growth rate on D T for a-axis and c-axis [20].
orientation close to the crystal growth direction in the undercooling method as well as in directional solidi®cation. In addition, Imagawa et al. [21] succeeded by controlling the growth direction in fabrication of the current lead with the Jc value of 7.0 104 A cmÿ2 (77 K, 0 T) by the conventional four-probe method. 5. DISPERSION OF THE FINE NONSUPERCONDUCTING PARTICLES
For application in high magnetic ®elds, the introduction of the pinning centers into the RE123 materials as the matrix is very important because the samples made by the processes of QMG, MPMG, etc. possess good properties in high magnetic ®elds. Therefore, understanding the mechanism for the introduction of the pinning centers into the RE123 is essential. In this section, size reduction of the particles as the pinning centers and introduction of these pinning centers are described. 5.1. Size reduction of nonsuperconducting particles Understanding of the coarsening behavior of the high-temperature stable phase particles in the liquid is extremely important since ®ner high-temperature stable phase particles are required to enhance magnetic ¯ux pinning as well as RE123 formation through the peritectic reaction. Izumi et al. [22] have investigated the coarsening rate of Y211 phase particles in the liquid by observations of pellet samples held at 10708C for dierent times. Figure 13 shows the holding-
Fig. 13. Holding time dependence of the mean radius of samples with and without Pt doping in a logarithmic plot [22].
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Fig. 14. Histogram of the Y211 distribution [27].
time dependence of the mean radii of Y211 phase particles in a logarithmic coordinate, assuming that the shape of the Y211 phase crystal is spherical, although the Y211 phase forms a rodlike shape crystal with some facets. The gradient of the lines is about the one-third, i.e. the mean radius increases proportionally to t1/3 (t, holding time). According to the above results, the mean radius of the Y211 phase particles increases proportionally to t1/3. Such coarsening behavior (one-third law) is often explained by the Ostwald ripening theory [23]. Izumi et al. [22] calculated the coarsening rate of the rod-like shape by modifying the theory. With several assumptions, the mean radius (R) of Y211 rods can be described as: R a DL Gt1=3 ;
5
where a is a constant, DL is the diusivity of Y in the liquid and G is the Gibbs±Thomson coef®cient. This equation has the same time dependence as that predicted by the Ostwald ripening theory and explains the experimental results of the Y211 coarsening behavior. The change in the coarsening rate is related to the change in the product DLG by Equation (5). Therefore, the eect of Pt doping may be explained by considering the fact that the dissolved solute changes the value of DLG which results in a change in the coarsening rate. By comparing the coarsening rates with and without Pt doping, it was found that the DLG value with platinum addition is about 10% of the value without Pt doping. As mentioned above, the size reduction of nonsuperconductor particles by doping an additive is eective. Recently, as the other examples, Kim et al. [24]
Fig. 15. Schematic drawings showing a particle in front of the solid±liquid interface and the necessary condition of particle pushing: (a) for Ds0>0, the force (Fi) due to the interfacial energy (Ds0) is conducive to pushing; (b) for Ds0<0, Fi is conducive to trapping. The drag force (Fd) is always conducive to trapping.
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Fig. 16. Schematically illustrated pushing/entrapment relationship between the critical growth rate (R*) and the critical radius (r*) of inclusion for dierent growth directions.
and Matsuoka et al. [25] presented that CeO2 is also eective as the additive for the Nd123 materials. As in the other methods, Kambara et al. [26] demonstrated that the size of nonsuperconductive particles is reduced by the Ba-rich compositions in the Nd123 materials. 5.2. Pushing/trapping for HTSC materials In this section, we present how to disperse ®ne high-temperature stable phase inclusions to the matrix of RE123. Figure 14 taken from Endo et al. [27] shows the histogram of the Y211 distribution for the Y123 crystal grown at two dierent undercooling, D T, values. From a quantitative analysis of the Y211 particles in the four regions with dierent D T and dierent growth directions, the number of Y211 particles in the D T = 30 K region is clearly much larger than that in the D T = 10 K region over a whole range of diameters. It should be noted that these phenomena are similar to a pushing/trapping behavior of foreign particles at an advancing solid±liquid interface during solidi®cation. Figure 15 is a schematic drawing showing a particle in front of the solid±liquid interface and the forces acting on the particle. There are two dominant forces: the drag force (Fd) due to viscous ¯ow around the particle, which moves together with the interface at R relative to the melt; and the force (Fi) due to the interfacial energy (Ds0). The other kinds of forces should also strictly be considered, such as gravity [28] and the van der Waals' force [29]. Furthermore, these forces should be dependent on the shape of the solid± liquid interface, which is determined by many factors, such as the thermal conductivity, pressure [30], distribution coecient [31], etc. To simplify the following discussion, however, we assume that the interface shape is planar while treating the two dominant forces. The necessary condition for particle pushing can be described as shown in Equation (6) [32]: Ds0 DsSP ÿ DsLP ÿ DsSL > 0;
6
where DsSP, DsLP and DsSL are the solid±particle, liquid±particle and solid±liquid interfacial energies, respectively. Fi is conductive to pushing if Ds0>0, while Fd is always conducive to trapping as shown in Fig. 15. According to the pushing/trapping theory, the critical size (r*) of a particle, larger than that trapped by the solid, is roughly determined by the critical growth rate (R*) and interfacial energy (Ds0): R ADs0 =Zr
7
where Z is the melt viscosity. Figure 16 shows the schematic relationship between the growth rate and radius of inclusion, which is derived using Equation (6) and considering R* r* = constant. From this ®gure, In the case of D T = 10 K, the critical radius (r*a) of a Y211 particle for the a-direction growth is smaller than r*c because of the anisotropy of R and Ds0. The parameters r*a and r*c at D T = 10 K are considered to be relatively large in the size distribution of Y211 dispersed in front of the interface because the total volume fraction of Y211 particles is much less than the calculated values from the lever rule using the initial composition for both directions. On changing D T from 10 to 30 K, the growth rates for both directions
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Fig. 17. Schematic illustration of pushing/trapping model for the Y123±Y211 system.
increase, leading to a signi®cant decrease in both r*a and r*c . Accordingly, even smaller Y211 particles, which were pushed at D T = 10 K, are trapped by Y123 crystals at D T = 30 K. The above mentioned phenomena are illustrated schematically in Fig. 17. In addition, Endo et al. [33] succeeded by controlling in fabrication of the single crystal bulk with the Jc value of 2 104 A cmÿ2 (77 K, 0.8 T) by magnetization. 6. HOMOGENEITY OF SINGLE CRYSTALS
To enhance the transport current, the larger volume of the RE123 as the matrix should be maintained in the conductor. In the case of the Y123 materials, the balance between the Y123 as the matrix and the Y211 as the pinning center is important because of a stoichiometric compound. However, in the cased of the other RE123 the additional problem of the solid solubility should be studied because of a nonstoichiometric compound. In this section, we describe the recent progress to minimize the solid solubility in the RE123 materials in contrast to the Y123 material. 6.1. Eect of oxygen partial pressure Nakamura et al. [34] reported the eect of oxygen partial pressure on the top-seeded solutiongrowth. The composition of the Nd123 single crystals grown in dierent oxygen partial pressure atmospheres was analyzed by ICP-AES and the results are shown in Table 1. From Table 1, it is recognized that with the same composition of solvent and solute, the deviation from the nominal composition (Nd:Ba:Cu = 1:2:3) is found to increase with increasing oxygen partial pressure atmosphere during crystal growth. This result indicates that a low oxygen partial pressure atmosphere during Nd123 crystal growth is found to be eective for minimizing the substitution of Nd ion into Ba sites. Figure 18 shows the temperature dependence of magnetization for the Nd123 single crystals shown in Table 1. It is obvious that it is possible to enhance the Tc value by minimizing the substitution of Nd ion into Ba sites, and in the research ®elds of the melt processing. Yoo et al. [35] found similar results. Recently, Yoshizumi et al. [36] reported the interesting results as the following. Figure 19 shows NdBCO quasi-ternary phase diagram in low pO2. This phase diagram diers from that in air shown in Fig. 20 [10], but resembles that of YBCO system. The solubility limit of Nd1 + xBa2 ÿ xCu3O7 ÿ d is very small (xmax=0.04). Accordingly, Nd123 superconductors made in the low oxygen partial pressure atmosphere are nearly stoichiometric, and this is the reason why the single crystals grown in the low pO2 atmosphere by Nakamura et al. show high a Tc of 96 K.
Table 1. Composition of Nd123 single crystals grown in dierent oxygen partial pressure atmospheres [34] p(O2) (atm) 0.01 0.21 1.00
Nd:Ba:Cu 1.01:1.97:3.00 1.07:1.95:3.00 1.10:1.90:3.00
K. KAKIMOTO and Y. SHIOHARA et al.
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Fig. 18. Temperature dependence of the normalized magnetic susceptibility for Nd123 single crystals grown in dierent oxygen partial pressure atmosphere [34].
6.2. Control of composition Recently, it is reported to minimize the substitution of RE ion into Ba sites by control of the composition. As an example, Fig. 21 by Yao et al. [37] shows the temperature dependence of magnetization of Nd123 single crystals grown from the liquid with dierent ratios of Ba/Cu. The eects of the Ba/Cu ratio in the liquid on Tc and superconducting transition behaviors are clear. With increase of the Ba/Cu ratio, Tc increases and the transition width D T becomes narrow, indicating that the ratio of Ba to Cu in the liquid plays a signi®cant role in controlling the Tc of Nd123 superconductor. From Fig. 4 (see Section 2), it can be seen that points on the liquidus through the tie lines connect to dierent points on the solid solution line. In other words, the substitution of Nd at the Ba site is a function of the Ba/Cu ratio in the liquid. With increase of Ba/Cu ratio in the liquid, the Nd substitution content decreases. When the Ba/Cu is
Fig. 19. NdO1.5±BaO±CuO quasi-ternary equilibrium phase diagram in 1% pO2 at about 1273 K [36].
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Fig. 20. NdO1.5±BaO±CuO quasi-ternary equilibrium phase diagram in air at about 1323 K [10].
approximately 0.8, Nd123 is close to the stoichiometric 123 composition. The Tc results are explained in accordance with this phase diagram. Also in the ®eld of other melt growth, Watanabe et al. [38] and Kambara et al. [26] reported similar results for the bulk materials. 6.3. Eect of growth temperature Kambara et al. [10] investigated the NdBCO quasi-ternary phase diagram for the region of dierent temperatures where Nd123 and liquid coexist in air. According to their results, it was found that the tie lines between Nd123 with the low substitution and the liquid phase lean toward the low substitution as the equilibrium temperature decreases. Therefore, it is considered that the Nd123 with the lower substitution is obtained from the liquid phase of the same composition under the lower temperature. As an example, Fig. 22 shows the relationship between
Fig. 21. Temperature dependence of magnetization for NdBCO single crystals grown in air with dierent ratios of Ba to Cu in the liquid [37].
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Fig. 22. The relationship between the growth temperature and x value in Nd1 + xBa2 ÿ xCu3Oy analyzed by EPMA [39]. The ¯ux composition of Ba/Cu = 3/5 was used.
the growth temperature and the substitution of Nd ion into Ba sites. This result is from Ishida et al. [39], in which the grown ®lm compositions are analyzed for the samples prepared by the LPE method with the ¯ux of Ba/Cu = 3/5. From the result, it is clear to minimize the substitution of Nd ion at the Ba site by preparation under a lower growth temperature, which agrees with the report by Kambara et al. [10]. 7. CONCLUSIONS
The enhancement of superconducting properties in the various methods for the crystal growth of the REBa2Cu3O7 ÿ d has been reviewed focusing on the recent progress in research. The study for the mechanism of the crystal growth of RE123 materials proceeded and the superconducting property also progressed. In the case of application to a conductor under a high magnetic ®eld, the introduction of pinning centers is very important and understanding the crystal growth mechanism becomes more eective. Recently, the Nd123 materials drew the attention of researchers because of its good properties in the high magnetic ®elds and their tremendous eorts have concentrated on the problem for solid solubility in the Nd123 materials. Furthermore, we believe that clari®cation of the growth mechanism of the RE123 materials will assist in further understanding of other oxide materials as well. AcknowledgementsÐThe authors wish to express their sincere gratitude to Dr. A. Endo for his achievement, to Dr. Kambara for making the eective phase diagram, to the other researchers in Division 4 of SRL-ISTEC for kindly providing their results. Parts of the studies presented in this paper were supported by the New Energy and Industrial Technology Development Organization for the R and D Industrial Science and Technology Frontier Program.
REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12.
J. G. Bednorz and K. A. MuÈller, Z. Phys. B 64, 189 (1986). M. Nakamura, Y. Yamada, T. Hirayama, Y. Ikuhara, Y. Shiohara and S. Tanaka, Phys. C 259, 295 (1996). B. J. Lee and D. N. Lee, J. Am. Ceram. Soc. 74, 78 (1991). B. J. Lee and D. N. Lee, J. Am. Ceram. Soc. 72, 314 (1989). K. W. Lay and G. M. Renlund, J. Am. Ceram. Soc. 73, 1208 (1990). T. B. Lindemer, J. F. Hunley, J. E. Gates Jr, A. L. Sutton, J. Brynestad, C. R. Hubbard and P. K. Gallagher, J. Am. Ceram. Soc. 72, 1775 (1989). S. I. Yoo and R. W. McCallum, Phys. C 210, 147 (1993). T. Wada, N. Suzuki, T. Maeda, A. Maeda, S. Uchida, K. Uchinokura and S. Tanaka, Appl. Phys. Lett. 52, 1989 (1988). M. Daeumling, J. M. Seuntjens and D. C. Larbalestier, Nature 346, 332 (1990). M. Kambara, M. Tagami, X. Yao, T. Umeda and Y. Shiohara, J. Am. Ceram. Soc. submitted, (1997). S. Jin, T. H. Tiefel, R. C. Sherwood, R. B. van Dover, M. E. Davis, G. W. Kammlott and R. A. Fastnacht, Phys. Rev. B 37, 7850 (1988). M. Murakami, M. Morita, K. Doi and K. Miyamoto, Jap. J. Appl. Phys. 28, 1189 (1989).
Melt growth processing of superconductive oxides 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39.
533
H. Fujimoto, M. Murakami, S. Gotoh, Y. Shiohara, N. Koshizuka and S. Tanaka, Adv. Supercond. 1, 285 (1989). Y. Yamada and Y. Shiohara, Phys. C 217, 182 (1994). T. Izumi, Y. Nakamura and Y. Shiohara, J. Cryst. Growth 128, 757±761 (1993). M. J. Cima, M. C. Flemings, A. M. Figuredo, M. Nakade, H. Ishii, H. D. Brody and J. S. Haggerty, J. Appl. Phys. 72, 179 (1992). N. Mori, H. Hata and K. Ogi, J. Jap. Inst. Met. 6, 648 (1992). T. Izumi and Y. Shiohara, J. Mater. Res. 7, 16 (1992). Y. Imagawa and Y. Shiohara, Phys. C 262, 243 (1996). A. Endo, H. S. Chauhan, Y. Nakamura and Y. Shiohara, J. Mater. Res. 11, 1114 (1996). Y. Imagawa, K. Kakimoto and Y. Shiohara, Phys. C 280, 245 (1997). T. Izumi, Y. Nakamura and Y. Shiohara, J. Mater. Res. 8, (1993). W. D. Kingery, H. K. Bowen and D. R. Uhlmann, Introduction to Ceramics, Wiley, New York, (1960). C. J. Kim, H. W. Park, K. B. Kim, K. W. Lee, I. H. Kuk and G. W. Hong, Mater. Lett. 29, 7 (1996). S. Matsuoka, M. Sumida, T. Umeda and Y. Shiohara, Phys. C 282±287, 531 (1997). M. Kambara, A. Endo, Y. Watanabe, Y. Shiohara and T. Umeda, International Workshop on Superconductivity. p. 202 (1997). A. Endo, H. S. Chauhan, T. Egi and Y. Shiohara, J. Mater. Res. 11, 795 (1996). S. G. Grigoryan, A. S. Oganesyan and A. G. Sarkisyan, Kristallogra®ya 28, 782 (1983). J. Potoschke and V. Rogge, J. Cryst. Growth 94, 726 (1989). G. F. Bolling and J. CisseÂ, J. Cryst. Growth 10, 726 (1971). A. A. Chernov, D. E. Temkin and A. M. Mel'nikova, Sov. Phys. Crystallogr. 21, 369 (1976). D. R. Uhlmann, B. Chalmers and K. A. Jackson, J. Appl. Phys. 35, 2986 (1964). A. Endo, H. S. Chauhan and Y. Shiohara, Phys. C 273, 107 (1996). M. Nakamura, H. Kutami and Y. Shiohara, Phys. C 260, 297 (1996). S. I. Yoo, N. Sakai, H. Takaichi, T. Higuchi and M. Murakami, Appl. Phys. Lett. 65, 633 (1994). M. Yoshizumi, M. Kambara, Y. Shiohara and T. Umeda, International Workshop on Superconductivity. p. 295 (1997). X. Yao, M. Kambara, T. Umeda and Y. Shiohara, Jap. J. Appl. Phys. 36, 400 (1997). Y. Watanabe, K. Miyake, A. Endo, K. Murata, Y. Shiohara and T. Umeda, Phys. C 280, 215 (1997). Y. Ishida, K. Kakimoto, Y. Shiohara and A. B. Sawaoka, Phys. C., (to be submitted).