Some recent developments in high-temperature superconducting oxides: a work review

185

Materials Chemistry and Physics, 34 (1993) 185-200

Invited Review

Some recent developments oxides: a work review

in high-temperature

superconducting

M. K. Wu Materials Science Center and Department

of Physics, rational Tsign Hua U~i~ers~~ Hsinchu (Taiwan, ROC)

(Received

November

September

4, 1992; accepted

9, 1992)

Abstract An overview of recent progress in high-temperature superconductivi~, particularfy the work carried out in our laboratory, is presented. The focus is on discussion of some of the problems and issues of current interest, especially those closely relating to practical applications of the high T, oxides. Typical experimental results on the development of the materials phase stabilization; crystal growth techniques; the effect of doping on the superconducting properties of the Y123 system; and the transport properties of the Y123 system in the presence of a strong magnetic field are presented and discussed.

1. Introduction The discovery of high-temperature super~onductivi~ in copper oxides [l, 21 has generated an unprecedented level of activity in laboratories around the world [3] during the last five years. The motivations are an interest in understanding the fundamental underlying mechanism and the potentially revolutionary applications of the high Tc materials on both small and large scales. Although new material systems have been found and their physical properties characterized, and many theoretical models have been proposed, no clear understanding of the mechanism responsible for high-temperature superconductivi~ has been established. Consequently, there are no science-based guidelines for researchers in their search for new high Tc materials. The picture that has emerged from the intense but scattered activity by the scientific and engineering community emphasizes the need for a systematic and interd~ciplina~ approach in order to understand these new superconductors. Most of the research activities carried out aim at developing our understanding of (a) the superconducting mechanism and its relation to structural and materials parameters; (b) the thermodynamic and kinetic factors controlling the stability of the appropriate phases; (c) the effects of processing parameters, thus enhancing our capability to control

0254-OS84/93/$24.00

and to synthesize reliable and reproducible materials; and (d) the novel superconducting properties and their implications for applications. It is generally accepted that the conventional BCS theory ]4] cannot adequately describe high-temperature superconductivity [5-71. A simple illustration of the difficulty encountered by the BCS theory is demonstrated by the near absence of the isotope effect [S]. Consequently, on this most fundamental of levels in the quest to understand the high T, superconducting phenomenon, the challenge is to find a novel pairing mechanism and/or condensation process which must together explain the observed high T,‘s and related phenomena. This challenge is being answered by groups taking different, but somewhat consistent, approaches [9-151. Some are based on modified BCS-type pairing, such as the spin bag approach [9]. Nevertheless, a preliminary assessment indicates that various phenomena or effects must be considered in the development of the theory. These include, to name a few: (1) The striking property of the normal state resistivities, which increase linearly with temperature in the Cu-0 systems [16]. (2) The intriguing temperature dependence of the Hall angle observed in the cuprate systems [17]. (3) A nonlinear thermoelectric power in the system, which implies the existence of strong electron-electron correlation [18].

0 1993 - Elsevier Sequoia. All rights reserved

186

(4) The unusual structural transitions [19, 201 from high-temperature tetragonal to orthorhombic and then to tetragonal symmetry at low temperature in the &.J3a0.&uQ-, system. A particularly intriguing and common characteristic of the high-temperature superconductor is its unusual phase diagram, as displayed schematically in Fig. 1, where the variance is directly related to the charge carrier concentration. In addition, a universal correlation between the transition temperature and the carrier densities has also been demonstrated from muon spin rotation [21] and Hall effect [22] experiments. A major barrier to the better understanding of this exciting new problem is the lack of relatively large perfect crystals. Existing single crystals of the high T, oxide compounds are either too small or infested with imperfections such as twin boundaries. Consequently, it is difficult to gain definitive info~ation for resolving the underlying mechanism responsible for high-temperature superconductivity. On the other hand, there are problems in achieving technological applications of this new class of materials. One difficulty that stands squarely in the way of application is achieving high current densities in high magnetic fields in bulk material. These problems may relate to the presence of grain boundaries, although there is some debate as to what aspect of the boundaries causes the deleterious effect. Some of the most plausible explanations include impurity segregation to boundaries and microcracking at boundaries attributable at least in part to the anisotropic coefficient of thermal expansion. It has been shown that fewer high-angle grain boundaries improve the current-carding capacity of bulk material 1231.Another, more fundamental difficulty is the complexity of the transport behavior in the presence of a magnetic field. There is also a lack of detailed understanding of such important issues as the effect of oxygen stoichiometry and ordering, and boundaries effects.

Fig. 1. Schematic phase diagram of the high-temperature superconducting oxides. The notations are P: paramagnetic state; AF: antiferromagnetic state; SC: superconducting state; and T and x are the temperature and carrier concentration, respectively.

In order to establish the structure-property relationships of the high T, materials and to achieve technological applications, it is necessary to develop processing techniques for materials displaying the desired characteristics. Depending on the particular quality considered, high-quality thin films or single-crystal materials are needed. It is also essential for optimum utilization that candidate materials be fully characterized with regard to electrical, magnetic and thermal properties. In the last five years, much progress in the development of high-temperature superconductivity has been made and the results have been well reviewed [3]. The main purpose of this article is to discuss some of the problems and issues related to the current development of high T, systems. Specifically, the topics will include problems related to the oxygen stoichiometry and the oxidation process; techniques for single-crystal growth; the effect of chemical doping on the superconducting properties; and the transport properties of the oxides in the presence of a magnetic field. This article is by no means a complete review; discussion will focus on those areas on which we have been working. Typical experimental results, particularly for the Y123 system, are presented and discussed. This article is divided into four sections. Sections 2 and 3 deal with the development of the materials phase stabilization and crystal growth techniques, Section 4 discusses the effect of doping on the superconducting properties of the Y123 system, and Section 5 presents some of our recent results on the transport properties of the Y123 system in the presence of a strong magnetic field.

2. Development of the materials phase stabilization The discovery of the first 90 K superconductor, YBa,Cu,O,_~ (Y123), was based on an empirical protocol [24] that related the structure and ionic size of the constituent species directly from the results of the La-M-&--0 (La214) (where M =Ba, Sr, Ca and Mg) system. This same argument was also used for the substitution of Bi for Y, which resulted in the observation of an unstable superconducting phase at 60 K in the Bi-Sr-Gu-0 system [25] and the discovery of Y-Sr-Cu-0 with T, at 80 K [26]. Recently, Xu and Guan [27] observed a correlation between T, and the ionic radii of the rare-earth ions in the Pr-doped REBa,CU,O, (RE123) system, where RE stands for the rare earth. They argued that the ‘dispersion’ of T, originates from the hybridization of the conduction electron wave functions with the localized Pr 4f states. These results strongly support the generally accepted picture that the presence of the Cu-0, planar structure is the most important ingredient in the formation of

187

high T, superconductivity in the cuprates. It was shown from the muon spin relaxation experiment that T, has a unique correlation [21] with the ratio of the superconducting carrier densities to the effective mass, i.e., n,/m*. T, reaches a peak and then decreases as n,/m* increases. The peak value of T, can be -40 K, = 90 K or = 125 K, depending on the particular local fine structure of the CuO, planes. The detailed mechanism responsible for this T,-structure correlation is not understood and is currently an extensively studied topic. It is well established that the superconducting characteristic of the oxide superconductors depends strongly on its oxygen concentration. For example, 90 K superconductivity exists only in a fully oxygenated Y123 system, the 40 K transition can be observed in La,CuO, only when x is larger than 4, and unusual structural modulation is observed in the 115 K Bi-based (Bi,Sr,Ca,Cu,O, or Bi-2223) cuprates [28]. It is believed that these observations are also strongly related to the Cu-0 plane local structure, as exemplified by the fact that the coordination number of the major Cu site is 6 for the 40 K system, 5 for the 90 K system and 4 for the 115 K system. Consequently, the relation of oxygen ordering to superconductivity has become one of the most important aspects of these materials and attracts special attention. For example, the effects of short- and long-range oxygen order on electrical and magnetic properties provide important clues for the understanding of the fundamental mechanism for superconductivity operating in these systems. Furthermore, an understanding of the effect of oxygen on the superconducting properties helps isolate the relevant structural features in these systems [29]. It is also obvious that an understanding of the oxidation process of the high T, oxides has technological importance. Therefore, much research has been devoted not only to obtaining a better understanding of the problems but also to discovering better and novel oxidation processes. Recently, the low-temperature electrochemical oxidation method [30] has attracted great attention owing to its simplicity and high efficiency. For instance, it was demonstrated that the nonsuperconducting La,CuO, compound can be made superconducting with T,=-50 K, as shown in Fig. 2, by this method [31]. Unfortunately, application of this process to the other oxide systems, such as Y123, seems nontrivial [31]. Another approach, which is more expensive and more sophisticated, is the high oxygen pressure technique [32]. This technique has become a must in many major laboratories because it provides a channel for the search for new high T,materials. The most appreciated example is the discovery of the infinite Cu-layer Ba-doped SrCuO, [33] system. On the other hand, by employing a low-temperature oxygenation process and maintaining an oxygen at-

: i,-.I c,,

5 ~:I-c‘;

1

_.. . .. .. ..

.

.-.-‘-~‘.7-.rl~‘~~,....*.~~ ,_.‘.._,.‘...,.,,,,lw..,._..._C C__,. -,.. ._.. I*_,.“ :

Fig. 2. Magnetic susceptibility as a function of the temperature of the La,CuO, after low-temperature electrochemical oxidation, showing the superconducting transition at =SO K [31].

0

1

Fig. 3. 2-V relationship temperature oxygenation

2

3

3

5

6

v (ClW at various temperatures of the lowof Y-Ba-Cu-0 (5-6-11) [34].

mosphere during electrical and magnetic measurements, Chen et al. [34] observed the thermally recyclable zeroresistance states with transition temperatures above 200 K in mixed-phase Y-Ba-Cu-0 materials. Although the high-temperature zero-resistance state was verified through careful measurements of the current-voltage characteristics, shown in Fig. 3, using multiple lead and contact arrangements, only diamagnetic-like deviations and hysteresis behavior at the same temperature as the resistive transitions were observed in the magnetic

188

measurements. To clarify the problem, we have also carried out a more detailed study [35] on the mixedphase Y-Ba-Cu-0 of 5-6-11 composition. The results of a.c. magnetic susceptibility measurements (with a frequency of 37 Hz and field < 1.5 G) show, in addition to a superconducting transition near 90 K corresponding to the Y123 phase, a much smaller diamagnetic-like anomaly with the onset near 270 K. Using the 90 K transition as a reference, it is estimated that the volume fraction of the high-temperature signal has an upper limit of 5%. The zero field cool and the field cool a.c. susceptibilities [35] of Sr-doped 5-6-11 samples clearly indicate that these samples also exhibit diamagnetic deviation with the onset temperature near 200 K. A hysteresis behavior is also present in this Sr-doped sample. In order to gain more insight into the origin of the observed high-temperature anomalies, we carried out an electron microscopy study. A typical transmission electron microscope image of a 5-6-11 sample exhibiting a high-temperature diamagnetic-like anomaly is displayed in Fig. 4. While the X-ray diffraction pattern shows a preponderance of the 123 phase, the TEM picture indicates the presence of a glassy-like phase.

Detailed structural and compositional analysis suggests that this glassy-like phase may be due to the intergrowth of the Y123 and Y,Ba,Cu,O,, (Y248) phases. A possible consequence is the presence of defect structures in the Cu-0 plane, which may lead to the anomalous behaviour. Although zero-resistance states under repeated thermal cycles were observed in the Y-Ba-Cu-0 (56-11) mixed phase by Chen et al. [34], we were not able to reprod.uce the same resistive transition in our study. On the other hand, stable diamagnetic-like anomalies are observed in both Sr-doped and undoped samples. The relatively small volume fraction of the magnetic anomaly may account for the absence of the resistive transitions. The above observations call strongly for a better understanding of the oxidation process and the development of better control of the oxygen stability in the high T, oxides. Many oxygen diffusion studies [36-381 have been carried out, especially in the Y123 system. These studies suggest that the out-diffusion is most likely surface-reaction-limited, while the in-diffusion is diffusion controlled. The diffusion mechanism is through the defects, which include the oxygen vacancies and twins. In general, several activation energies have been observed, corresponding to the out-diffusion and the in-diffusion (depending on the oxygen concentration). The higher activation energy for the out-diffusion process suggests that the surface barrier, which can be affected by the grain boundaries and impurities, plays an important role in controlling the oxygen concentration. If a good oxygen catalyst is used in contact with the oxide, then better stabilization of the phase might be achieved through the limiting oxygen outdiffusion process. A recent study on the Ag-added Y123 composites [39] indicates that oxygen stabilization is achievable using Ag as a control agent. Recently, we found that the addition of a small amount of a transition metal element such as MO, substituted for Cu, can stabilize the high T, 123 phase in the Y-Sr-Cu-0 system (details of the results will be presented in Section 4). However, these results suggest the important concept of establishing the phase stability in the high T, oxides through the introduction of either pure metals or metal oxides.

3. Development

Fig. 4. TEM image a glassy-like phase.

of a Y-Ba-Cu-0

(S-6-11)

sample,

showing

of single-crystal

growth techniques

Although several high T, oxide superconductors have been successfully synthesized, the amount of thermodynamic information accumulated for these systems is still rather limited. Further progress in the development of new and existing oxide superconductors depends crucially on our having sound thermodynamic information and data. In particular, equilibrium phase dia-

189

grams, or portions of them, characterizing solid-liquid as well as solid-solid reactions are needed in order to be able to control fabrication and heat treatment processes. In the early development of high T, superconductivity, it was found that a nonequilibrium processing technique that requires a high processing temperature and fast quenching of the nearly molten oxides to room temperature was very effective for producing new high T, materials. For instance, the observation of superconductivity in the first 90 K Y-Ba-Cu-0 superconductor and the subsequent discovery of Y-Sr-Cu-0 and Bi-Sr-Cu-0 [25,26] all occurred through this approach. Although the detailed kinetics of the formation of these high T, oxides remains unknown, the results suggest that the nonequilibrium process is more viable for the formation of new high-temperature superconducting oxide compounds with appropriate fine and gross microstructures. Furthermore, recent developments in the processing techniques for the preparation of bulk materials with enhanced superconducting characteristics [40], e.g., flux pinning strength have also employed a similar nonequilibrium process as the first step. The lack of detailed knowledge of the equilibrium phase diagram for high T, oxides also hinders progress in the growing of high-quality, large single crystals. Most of the current crystals are grown [41] from the flux based on a trial-and-error approach to determining the optimum composition. Recently, a liquid-phase processing method [42,43] has been developed to grow YBa,Cu,O,, crystals from a YBa,Cu,OJAg,O composite material. Single crystals can easily be extracted from a specimen that has been sintered at 1005 “C for more than 8 h. The details of the crystal growth processes are as follows. Y123 powder prepared by solid state reaction was mixed with Ag,O powder in a weight ratio of 1:3. The powder was mixed and pressed in the form of a disc (1 mm thick, 10 mm in diameter) under 3 ton cmP2 pressure. Each pellet was laid flat on top of a piece of gold foil in an aluminum crucible and then introduced into a tube furnace with a continuous flow of oxygen. The specimens were initially sintered at 950 “C for 6 h, then slow-heated to 1005 “C. Exaggerated grain growth of Y123 occurred after 8 h at this temperature. The specimens were cooled at a rate of 6 “C h-’ to 950 “C and held there for 6 h. They were then slowly cooled to 500 “C and held there for 20 h for oxygen uptake before being cooled to room temperature. During the sintering process at 950 “C, the growth of small Y123 facet-like grains was observed on the surface of the pellet. The Ag particles resulting from the dissociation of Ag20 scatter primarily in the void, and cause the pellet to become denser. Silver beads

began to form on the surface of the pellet as the specimen was heated to a temperature above 980 “C. No conspicuous grain growth could be seen in the initial few hours of sintering at 1005 “C. EDX analysis of the whole specimen shows that a small amount of Y,BaCuO, phase and Ag-rich particles dispersed in the grain or at grain boundaries. Ba-Cu-0 flux, resulting from partial melting of Y123, was extruded out of the grain and scattered at voids. Most of the Ag particles embedded in the flux react with BaCuO and form a complex network structure. As the sintering time at 1005 “C increased to more than 8 h, large stacked plate-like Y123 crystals appeared. Crystals are arbitrarily distributed on the surface of the pellets. The top surfaces of the crystals are full of hillocks or waves, as shown in Fig. 5. EDX analysis indicates that most hillocks are Y123 phase with an excess of Y, while some of them are Y-rich or Ag-rich precipitates and there appears to be a composition gradient around the precipitates. The Ag and Ba-Cu-0 melts are distributed at grain boundaries and grain corners. Some of the Ag-rich small liquid particles and Y,BaCuO, precipitates were trapped in the grain or in the facet grain boundaries because they were unable to diffuse out in time, while plate-like grains rapidly contact and coalesce. Most of the Ag and Ba-Cu-0 melts percolate through the compact and react with the gold foil underneath. Except for a small amount of Ba-Cu-0 melt that remains in the void, Ba-Cu-0 melts were jammed to the sample surface and reacted with Y,BaCuO, phase and Ag-rich particles, which act as nucleation and growth centers, resulting in the hillock microstructure. Wellstacked plate-like grains are arranged in the a-b direction and tend to increase in size by further coalescence and growth. Single crystals of 3 X 2 X 1 mm in

Fig. 5. Optical micrograpb of the crystals grown from the Y123/ Ag composite. The surface is full of hillocks, which contain Yrich Y123 particles, according to EDX analysis [42].

190

size can be grown using this process with more than 16 h of sintering at 1005 “C. From the microstructural and chemical composition analysis, we propose the following grain growth mechanism. The capillary force from a wetting Ag liquid introduces a stress at the Y123 intergranular contact point. This stress causes preferential decomposition of the Y123 solid at the contact points. A fraction of the Y123 solid is decomposed into Ba-Cu-O-rich liquid and Y,BaCuO, precipitate. The dissolution of Ag and the Ba-Cu-0 melts open up the pores and form egresses through which the liquid can percolate through the compact. Y123 grains are brought into contact by a wetting liquid after liquid flow. Contacting grains of dissimilar sizes fuse into a single grain by a continuous process of directional grain growth and grain reshaping. At this stage, the mechanism of shape accommodation of Y123 involves the dissolution of small grains and reprecipitation on large grains. The percolating behavior of silver liquid causes the volume fraction of the liquid to decrease, which in turn causes the degree of shape accommodation to increase. Y123 crystals prefer to grow along the a-b plane and then form plate-like grains. New contacts between grains are further induced by gravity, thermal motion and settling. A low misorientation angle between the contact grains results in a high probability of subsequent coalescence. With a high contiguity, there is more grain contact and more opportunity for coalescence. Y123 crystals grown using this technique are heavily twinned. They show zero resistance at 92.2 K with a width of less than 0.4 K in zero field. The superconducting resistive transition in magnetic fields is broadened with zero resistance and decreases to 70 K in a field of 15 tesla, as shown in Fig. 6. Large domains of l-2-3 consist of stacked, low, crystallographically misoriented grains. By magnetization measurements, it was found that J, is about 2~ lo4 A cm-’ at 77 K and 1 tesla. Another novel approach developed to grow singlecrystal oxide superconductors is the low-temperature anodic electrocrystallization method in a potassium hydroxide molten flux [44]. This method was first demonstrated to be very effective in growing Ba,_,KBiO, (BKBO), with various values of x. Recently, it was also shown that values of x greater than 0.5 [45], which was claimed to be the solubility limit of this compound system [46], can be obtained only by this crystal growth technique. For example, as displayed in Fig. 7, the x = 0.56 bismuthate crystal shows = 8 K superconductivity. Transport measurements on these crystals (Fig. 8) show that the temperature dependence of the normal state resistivity follows a simple power law up to the second order, while the Hall coefficient is temperature independent. Based on these transport measurements,

0.00006

h -s 0.00004 h .% .$ .I? t (I: 0.00002

0.00000 70

75

60

65

Temperature

90

95

I 0

(K)

Fig. 6. Resistivity of Y123 crystal growth from Y123/Ag composite under various magnetic fields up to 15 tesla [43].

-

5’oE-007 E

3

2.

-1.3E-021:

$ 3 Fz -5.OE-007: ; 2 ti 5 -l.OE-006: m c!.' , -1.5E-006:

-2.0E-006i""""~""""'I~"""~~'~~~""""""""'~~"""~ 0 5 10 15 20 TEMPERATURE

Fig. 7. A.C. magnetic susceptibility K [45].

25

:

(I<)

of B%,,K,,,,BiOX crystal, T, = 8

the bismuthate BKBO does not exhibit the l/T Hall anomaly [45] characteristic of the cuprates. This difference suggests that the Hall anomaly is not a universal feature of high T, oxide superconductors and is more likely a consequence of the layered structure of the cuprates. In contrast to the thermal reaction methods, the lowtemperature electrochemical method can provide several important advantages for the synthesis of bismuthates. First, it avoids the product nonstoichiometry caused by the volatility of alkaline oxide at high temperatures. Second, the oxidative power is readily delivered to the desired crystal/solution interface, making precise control of the oxidation possible. Furthermore, large, uniform crystals can be obtained. For instance, a 36 h experiment produced crystals with 2 mmx 2 mm cubic facets. The exact crystal growth mechanism of the low-temperature anodic electrochemical tech-

191

Temperature

(K)

Fig. 8. Temperature dependence of resistivity and Hall coefficient for BKBO crystals [45].

nique in the bismuthate system is not completely clear. The temperature, applied potential, and flux composition are correlated with product properties. To obtain crystals of the correct stoichiometry, precise control of the experimental conditions is necessary. A preliminary study has shown that a similar technique can also be used to grow La,CuO, crystals [31]. The development of techniques for the growth of high-quality single crystals not only provides the desired materials for fundamental research, but also provides important guidelines for the development of novel approaches to fabricate bulk materials for practical applications. The process using the Ag-added composite for large grain growth has been used as the basis of the continuous process for bulk texture processing, while the low-temperature anodic electrocrystallization technique has been applied as a novel process for lowtemperature oxidation in various high T, oxides.

4. S-doping

of the Y123 system

In order to gain more insight into the mechanism responsible for the occurrence of superconductivity in the oxide systems, studies using ion substitution in various cuprate superconductors have been carried out. Taking the Y123 system as an example, substitution of the Y ion by all the rare earth elements revealed that except for Ce, Pr and Tb ions, the RE123s remain superconducting at 90 K [47]. The substitution of Cu by most of the metal ions [48] resulted in degradation of the superconductivity, though the magnitude of the suppression depends on the particular ion species. In general, greater suppression occurs when the ion sub-

stitutes for the square planar Cu and is less pronounced when the substitution is at the linear chain Cu site. Though the partial substitution of the Ba ion by other divalent ions resulted in a relatively minor effect on the superconducting transition temperature, most of the early attempts at complete substitution failed to maintain the crystal structure [49]. These results strongly suggest that the Cu-0 square planar structure is the most important component in the superconductivity of this oxide system. Recent studies of the isostructural (Y,_,Pr,)123 series [27, 501 also support the same conclusion. However, there were two debatable mechanisms proposed for the reduction of superconductivity in the Pr-doped system. One is based on the pairbreaking mechanism due to the presence of magnetic impurities; the other is attributed to the reduction of hole carriers by doping. The current data cannot clearly resolve these two different pictures. Further detailed study along this line has become one of the current interests. It is noted, based on our recent results [51], that a more careful treatment of the material, especially the elimination of the phase separation problem, is needed before a clear resolution of this issue can be achieved. Recently the complete substitution of the Ba ion by Sr has successfully been made while maintaining the 123 structure. These results have consequently generated an enormous amount of activity. The earliest report suggested the possible existence of a YSr,Cu,O, (YSCO) phase with T, about 80 K in a multiphase system [26]. Later reports [52-541, including one using a high-pressure synthesis route [52], also did not provide samples with a clean superconducting phase. Cava et al. [55] found that by including Pb, there exists a series of new compounds Pb,RESr,Cu,O, with T, = 40 K. It has tetragonal crystal symmetry and the structure is similar to that of Y123, but there is an unusual Pb-0 cage [56] embedded in between the Cu-0 planes. This discovery led to the successful synthesis of a new series of cuprates, the so-called 1212 system, which includes the Tl-based Tl-Y-Sr-Ca-Cu-0 compounds. This system has attracted great attention because of its relatively high T, and because of its superconducting characteristics, which are comparable to those of Y123 in magnetic fields. It has also subsequently generated many new superconducting systems, such as the Pb-Hg-Sr-Ca-Cu-0 [57] system discovered very recently. On the other hand, by introducing a small amount of MO to partially replace the Cu atom, single-phase YSCO was successfully prepared with T, around 40 K [58]. The presence of MO was found to reduce the reaction temperature and is essential for the formation of the high T, YSCO phase. Similar phase stabilization was also made by the introduction of Fe, Al, Co, etc. [59]. These observations provide us with an opportunity

192

to study the role of the Ba ion in the superconductivity of Y123. Therefore, we have carried out detailed studies of the Y(Ba~_~Sr~)~(~~_*.~Mo~.~)O~ system, with x ranging from 0 to 1 [60]. The reason for varying the MO and Sr contents simultaneously is to maintain the phase purity of the whole system. Depending on the strontium concentration, the superconducting transition temperature varies from 90 to 40 K. The results of resistivity measurements show that T, decreases monotonically from 90 K for x==0 to 40 K for x = 1. The normal state resistivity above T, increases with increasing Sr concentration, which may be due to the combined effects of increasing Ba/Sr disorder scattering or decreasing oxygen content from the increase in MO content [60]. The normal state resistivities of the samples clearly exhibit a gradual change from linear to nonlinear temperature dependence as x increases. As x increases, linear temperature dependence of the resistivity is still observed at high temperature, but a downward curvature starts to devefop at temperatures far above the critical temperature. When x approaches 0.4, the downward curve begins even at 300 K. However, for x>O.4, a slight upturn of the resistivity gradually appears at temperatures near the transition. In general, the resistivity in the whole range can be described by a simple power law. If we fit the resistivity to the power law T’, the best fit for x < 0.4 gives z < 1, and for x > 0.8, z > 1. The above observations are qualitatively consistent with the proposed model based on the existence of van Hove singularities near the Fermi level [1.5]. The detailed X-ray diffraction patterns of the Modoped Y-Sr-Cu-0 are shown in Fig. 9. A tetragonal structure with lattice constants a = 3.818 A andc = 11.555 A was determined by Rietveld refinement [58]. The atomic positions of space group P4fmmm and the iso-

( A)--Y-Sr-Cu-o 03)---Yz03 (C)-YzSrQ,, (Djdhesubtracted

tropic Debye-Waller factors listed in Table 1 were determined using a program based on the stoichiometry of YSr&u&&+y7 where the values of the Debye-Wailer factors were confined to between 0.3 and 4.0 (the corresponding atomic rms displacements are 0.05 8, and 0.25 A). The final R factor is 12.6%. Atoms O(3) and Cu(1) on the basal plane can be considered to be relatively unstable owing to their large Debye-Waller factors. C&,1) atoms were dragged along by the unstable O(3). The distance between the 010, planes is 3.305 A, which is smaller than that for YBa2Cu306+,,, 3.388 A. Table 2 gives the respective cation-oxygen bond lengths. The results are not much different from those for the Y123 compound. Table 3 displays the lattice parameters and the unit cell volume V from for X-ray analysis the compound series Y(Ba, _,Sr,),(Cu,_O,,Mo,,,)O~. The lattice constants decrease with increasing Sr content. A phase transition from orthorhombic to tetragonal structure at x=0.4 is clearly shown, This transition coincides with the sudden drop in T, for n between 0.4 and 1. The observed structural phase transition is most likely due to the simultaneous increase in MO with Sr content. As shown previously [60], MO ion is in the “r 6 valence state and sits in the Cu(1) site. When the MO concentration reaches a critical value, extra oxygen will be adsorbed and the structure changes from orthoTABLE 1. Position parameters and isotropic Debye-Waller factor for the tetragonal structure of YSrz(CuzsMo,,a Oy based on space group P#mmm, a=3.818a and c=l1.555 B Atom

Y Sr Cul cu 2 01 02 03

Wyckoff notation

x

Id 2h la 2g 2g 4i 2f

0.5 0.5 0 0 0 0 0

h

20

40

GO

%I

28(dcgroZ) Fig. 9. X-ray diffraction patterns and (d) Y&, (c) Y2SrO4 YSr&%&%a)4 iS81.

z

0.5 0.5 0 0 0 0.5 0.5

0.5 0.19 0 0.36 0.16 0.36 0

B

0.04 3.86 2.53 2.88 6.35 3.01 0.78

occu- Madelung panty

potential

1 1 1 1 1 1 0.47

- 82.74 -49.31 - 138.18 - 56.04 - 29.24 -41.33 - 29.67

pattern TABLE

0

y

of (a) YSr&Zuz,MoO,&,, the subtracted pattern

(b) of

2. Cation-ovgen

bond lengths of YSraCuz.BMoo,,O,,

Bond

Length (A)

Sr-O(l) Sr-O(2) ?&o(3)

2.712 2.763 2.884

Y-O(2)

2.471

Cu( l)-O( 1) Cu(l)-o(3)

1.842 1.904

Cu(2)--w) G(2)-O(2)

2.273 1.905

193 TABLE 3. Superconductivity transition temperature TC,,lattice parameters a, b, c, unit cell volume V and b/a ratio for the series Y(Bal-,Sr~)z(Cu,-,,,Mo,.,)0, u

b

C

:I

(4

(4

(4

&I

92 90 85 74 68 61 55 48 42

3.821 3.818 3.821 3.838 3.831 3.827 3.808 3.813 3.810

3.889 3.879 3.857 3.848 3.831 3.827 3.808 3.813 3.810

11.691 11.650 11.644 11.618 11.595 11.607 11.544 11.556 11.548

173.726 172.550 171.646 171.610 170.145 169.992 167.367 168.044 167.589

x

0.0 0.1 0.3 0.4 0.5 0.6 0.8 0.9 1.0

b/a

1.018 1.016 1.010 1.002 1.000 1.000 1.000 1.000 1.000

rhombic to tetragonal. The possible causes of the large Tc suppression when the system becomes tetragonal include: compression of the lattice; the change in the carrier concentration; and the structural disorder. The first possibility is in contrast to the results for the effect of pressure on Y123 [61]. From the Raman scattering measurements, we found that the 330 and 500 cm-l modes have the most prominent peak shifts as x increases. For x=0.9, the 330 cn-’ mode shifts by 30 cm-’ toward lower frequencies while the 500 cm-l mode shifts by 30 cm-’ toward higher frequencies. The 440 cm- ’ mode also seems to shift to a higher frequency as x increases, but the fluctuation is large. Furthermore, there is a hump at 620 cm-‘, which becomes more pronounced when x is greater than 0.6. The 500 cm-’ mode shift to a higher frequenq as x increases can be understood as reduction of the c axis owing to the replacement of Ba with Sr. Consequently, the Cu(l)-O(4)-Cu(2) bond distance is reduced. Yamamoto et al. [62] made the lattice dynamics calculation for YBa,_,Sr$u,O, based on the shell model. Their main result is that both the 330 and 500 cm-’ modes increase as x increases. However, both their and our Raman data show that the 330 cm-’ mode has the opposite behavior. The explanation given by Matsuda et al. [63] is that the 330 cm-’ mode shift to lower frequency is due to the Y-(0(2), O(3)) bonding becoming weaker. This suggestion is not consistent with the data for the volume change with the Sr doping. An alternative and more likely explanation is that because Sr is smaller than Ba, the O(2) and O(3) atoms can be attracted to Sr. Therefore, although the a and b axes decrease as x increases, the O(2)-Cu(2) and O(3)-Cu(2) bond lengths may increase. This possibility is consistent with the orthorhombic-to-tetragonal phase change. Nevertheless, the observed change in the vibration mode cannot be used to explain the T, change of the series, especially the sudden change at x=0.4.

One unusual observation in our measurements is the presence of an anomalous peak at around 620 cm-‘. Hangyo et al. [64] observed a similar peak at 610 cm-’ in YBa,(Cu,_,Ni,.),O,_, for x greater than 0.1. They attributed the peak to the presence of Y,BaCuO, (green phase). Morioka et al. also found a similar hump at 575 cm-l, which was associated with the oxygen disorder, in YBa,(Cu,_,Co,),O,_, [65]. The X-ray diffraction data of our samples show no sign of the green phase. It is most likely that the broad mode in our observations originates from the induced crystal symmetry change through the introduction of the O(5) atoms. The vibrational modes of the O(1) oxygen which are infraredactive and Raman-inactive [66] in orthorhombic Y123, may be affected strongly by the disorder of O(1) oxygen. Therefore, the observed 575 cm-l Raman band is from the O(l)-Cu(1) stretching mode, which becomes Raman active owing to the disorder of O(1) in the tetragonal phase or the introduction of the O(5) atoms. The concentration dependence of this mode is similar to that of the T, variation. This leads us to propose that the T, depression is mainly from the electronic structure change due to the carrier density change. Indeed, recent Hall measurements [67] on this system confirm this suggestion. Therefore, we conclude that in the Sr-doped Y123, the change in the Cu-0 chain plays an important role in determining the change in Tc.

5. The resistive magnetic field

transition

of the Y123 system in a

Experimental results, such as the evidence for intragrain Josephson junctions and the development of a tail in the resistive transition in the presence of even moderate applied fields, indicate that the phenomenology of the high T, oxides is significantly different from that of the classical superconductors [68]. These observations have given rise to an unprecedented amount of activity in the search for a better understanding of the transport characteristics of the oxides and the effort to improve the material properties. A rationale for the understanding of these behaviors has been proposed based on the intrinsically short coherence length [69]. It is believed that in bulk high-temperature oxide superconductors, the critical current density (J=) and the critical magnetic field (H=) are controlled by details in the crystal structure, the microstructure, and the defect structure, including the phase boundary structure. In the last five years tremendous progress has been made concerning the problems related to the critical current densities of these materials. Specifically, the grain orientation has been shown to affect the critical current by about an order of magnitude. Material with

194

critical currents near the theoretical limit was obtained in properly oriented thin films [70]. It was demonstrated that the addition of AgO to form the RE-123/AgO [71] composites resulted in unusual magnetic properties and enhanced critical current densities. A decrease in the ratio between magnetization-determined .I, and transport-measured J, was also found in these composites [72]. Work by Tien ef al. [73] also revealed that the AgO or Ag is a mechanical processing aide. Not only is hot isostatically pressed (HIP) Y123/AgO [73-751 100% dense, but the bulk shapes were also relatively free of microcracks. Hot compression tests also showed that the composite superconductors are more hot-workable. It is noted that even without AgO, significant hot deformation (without cracking) of Y123 has been achieved in the appropriate temperature and strain rate ranges. Noble metal composites of RE-123 and monolithic RE-123 have been used to examine the basic features of J, - separating the flux pinning effects from the weak-link problems at grain boundaries and prior particle boundaries from other intragrain currentaffecting features such as dislocations and twin boundaries. An important issue being considered is the correlation between the coherence length, the grain boundaries and the electronic properties. More specifically, for the Ag-added composite, it is important to consider how thick the residual Ag should be between grain boundaries and how the choice of RE ion affects this thickness. It was shown that with proper heat treatment, silver increases the amount of flux pinning substantially. Silver also appears to wet the surfaces of the RE-123 well, segregating to prior particle boundaries and grain boundaries. It is believed that the Ag on boundaries enhances oxygen diffusion. This may be a necessary feature for bulk material, since preliminary data have shown oxygen diffusion to be slow in dense oxide superconductors. The presence of strong flux pinning indicates that silver might be introduced at finer sites such as dislocations, twin boundaries or precipitates. Nevertheless, the present improvements are still below the requirements for practical applications, especially in the materials with bulk form. The most detrimental effect is the rapid suppression of the critical current densities in the presence of a magnetic field. The lack of detailed knowledge of the dynamics of the flux lines of the high T, materials has strongly hindered progress. Muller et al. [76] first demonstrated the existence of the magnetic irreversibility line in the fieldtemperature (H-T) phase diagram in the ceramic samples. This led them to propose the importance of the glassy structure in the high T, materials. However, a similar effect observed in single crystals has led others to suggest the significance of the thermally activated flux creep [77]. These results have subsequently gen-

erated numerous experimental [78-831 and theoretical [84-881 attempts to understand the nature of the flux motion in the high T, oxides. Currently, the irreversibility line is characterized in two main ways - either as a phase transition of the flux line or as a region where the flux creep is enhanced in the new superconductors that possess the characteristics of short coherence length, long penetration depth, large anisotropy and high transition temperature. We recently carried out a series of detailed I-V characteristic measurements in fields [89, 901 on the Y 123 single crystal and thin films of different thicknesses. Our results strongly support the picture of a vortex-solid (or glass) to vortex-liquid transition predicted by the phase transition model [86] based on the presence of random disorder. However, the results also suggest that the detailed scaling behavior near the transition is dictated by the sample defects. Typical temperature dependence of the resistance in a magnetic field for the crystal and thin film is shown in Figs. 10 and 11, respectively. The zero-field transition width is within 1 K for the crystals, and about 1.5 K for the thin films. Differences in the field broadening of the resistive transition for the samples were observed. Particularly, the distinct ‘knee’ in increasing field is observed in the crystal and thicker films, but not in films with a thickness of less than 2000 A. Figure 12 shows the I-V curves at 4.6 T for Y123 crystal and a Y 123 thin film ( = 1 pm thick). Clearly, for both samples there exists a temperature (T,) at which the slope of the log( I+log(l) curve is exactly linear. In fact, we have observed similar behavior in all the samples studied. For temperatures below Tg, a downward curvature is evident at all currents, while above T,, the curvature is upward and becomes linear when the temperature is above the temperature where the ‘knee’ appears in the R-Tcurve at the same field. We have also determined ZE-004 ***** O.OOOT o*ao* 1.164T ~&~~*2.326T ===o 3.492T q

OE+OOO 70

75

80

85

90

95

100

Temperature (K)

Fig. IO. Temperature dependence of the resistivity of Y123 crystal (grown using the conventional flux method) under different magnetic fields up to 9 tesla [88].

195

E

3I

70

75

80 Temperature

85

90

95

I

(I()

Fig. 11. Temperature dependence of the resistance Y123 thin film under various magnetic fields.

of Ag-doped

the true T, (R=O) where the resistance is zero within the resolution of the measurement. T, and T, at different magnetic fields so obtained are plotted in Fig. 13; the results clearly show that these temperatures are apparently identical. The results suggest that the data can be described in the context of the vortex-solid to vortex-liquid transition picture [86]. To make the analysis more general, we fit our data at T, to a power law, Ea./“. The derived exponents are listed in Table 4 for the crystal and Agdoped thin film. Figure 14 shows the derived exponent cy for different samples. The results indicate that the exponent (Y depends on both the magnetic field and the sample thickness. For thicker samples, (Yfirst decreases in low fields and then increases until the field reaches a critical value, after which the exponent remains constant up to the highest field. On the other hand, for thin samples (=: 1 pm thick), the exponent remains constant throughout the whole range of the magnetic field. Furthermore, it is confined between the values for the Y123 thin film and those for the Y123/Ag thin film. Since Y123/Ag film is expected to contain more complicated defect structures, the observed difference in (Y, which can be viewed as a characteristic of the pinning strength, is likely due to the different interaction between the flux and the pinning centers. This observation suggests that there possibly exists a crossover of the flux motion from 3D to 2D. At high field the pinning center density is comparable to or even less than the flux density. The interaction between the flux lines becomes more important. This effect will suppress the possibility of flux entanglement. Therefore, the flux lines move like rigid bars and the motion is more 2D-like. By considering a distribution of the pinning potential, the critical exponent may vary but the trend of field dependence will be maintained. Depending on the film thickness, the critical field at

Y123 B=1.6

10

(b)

-I

th,n f,lm T?Sl,,

10

-J

Current

10

-z

(A)

Fig. 12. I-V curves of (a) Y123 crystal at H=4.6 tesla, where the temperature varies from 85 to 78 K, (b) a Y123 film at H=4.6 tesla, where the temperature ranges from 79 to 76 K, the temperature difference between successive curves is 0.3 K 1891.

which the dimensional crossover occurs varies. For example, the critical field for a 8000 8, thick film is about 0.5 T. It is expected that the same dimensional crossover will also occur in the single-crystal sample when the applied field is large enough. If we simply use the sample thickness for scaling, the crossover field for the single-crystal sample ( = 20 pm thick) is estimated to be at c. 12 T. The theory also predicts that the resistivity as a function of the current density satisfies a scaling relation (E/J) = tYE(Jt -p/T), where t is the reduced temperature, I (T-T&T, I. A typical result of the data-fitting using the scaling relation is shown in Fig. 15. The exponents p and y, extracted from the data-fitting, for the crystal

196

0

74

76

80

78

82

84

Temperature

Fig. 13. Field dependence

86

88

90

(K)

of Tg and T, (R=0) of Y123 crystal

WI-

2

0

92

4

6

8

10

B (Tesla)

Fig. 14. The critical exponents of the Y123 crystal and thin films as a function of magnetic field [89].

TABLE 4. Tr and critical exponents a, fl, y of Y123 single crystal and Ag-doped Y123 fdms at different 3 fiefds Ag-doped

Crystal (= 20 pm thick) (Y

P

Y123 film (Y

Y

B

Y

&

-

86.0

0

-

-

-

0.80

0.90

84.9

0.58

1.5

3.9

2.5

2.05

0.70

0.80

83.6

1.16

1.6

3.6

2.5

2.91

1.94

0.75

0.75

84.4

1.75

1.7

3.5

2.8

85.3

3.49

2.04

0.85

0.90

80.8

2.33

1.8

3.5

2.8

84.2

4.07

2.07

0.65

0.70

79.5

2.91

1.8

3.4

2.8

83.4

4.66

2.11

0.65

0.70

78.2

3.49

1.8

3.4

2.8

82.2

5.24

2.07

0.65

0.70

77.0

4.07

1.8

3.4

2.9

81.3

5.82

2.20

0.60

0.75

75.5

4.66

1.8

3.3

3.0

80.5

6.40

2.19

0.70

0.85

73.5

5.82

1.9

3.2

3.0

79.3

6.98

2.36

0.70

1.00

70.9

6.98

1.9

3.3

2.9

78.3

7.57

2.61

0.70

1.20

69.0

8.15

1.8

3.3

3.0

77.3

8.15

2.73

0.70

1.20

“0

93.3

0

-

-

88.5

1.75

2.08

87.4

2.33

85.9

c:

i3

k

2,

t

[O

and Ag-doped Y123 thin film are also listed in Table 4. The exponents satisfy a universal relation y= p(fr - l), as clearly shown in Fig. 16 for the crystal sample. This result is consistent with the prediction by Fisher et aE. [86], provided that the dimension is generalized to include the case of 6#3. Figure 17 displays the H-Tcurves for various samples, including a Ag-doped Y123 thin film. For the single crystal and Ag-doped Y123 thin film ( = 4500 A), the irreversibility line can be fit to the functional form of Ha[l - Z’,(Z3’)ITg(0)]n, with 17-3f2. For the thin film

-.L._.‘-...__L-._.&~~10

10 =

’

10

10 ’

l

.i..

J 1-P

Fig. 15. Data-fitting of the I-V curves of Y123 crystal using the scaling relation [88].

0

4

2

B-FIELD

6

A 6

(T)

Fig. f6. The field-dependent critical exponents CY,/?, y of Y123 crystal, showing the universal relation y= P(a- 1).

without Ag doping and = 1 ym thick, q is =5/3. For much thinner samples, the data do not follow a simple power law. A comparison of these results with the fielddependent exponent a; displayed in Fig. 14, shows that

197

.

u

.

I

‘1

.

. . :

. .

.,I

_. .

* i

.n

.

CI ‘7

n

.

0

*.A “G

-1--~-11

L-Yli

0.6

ilLu-ll

0.7

0.6

0.9

T,(B)/T,(O)

Fig. 17. H-T curves for various samples. A curve with smaller slope implies that the sample has a weaker pinning strength.

6EF004

-

-5E-004

u

YBCO (Tz,=90 5K) YBCC,‘Ag (ic,=90

5K)

i “E

4E-004

2

/

f

1 EGaIa

OE+OOO

. 50

100

150

2

TEMPERATURE(Kelvin) Fig. 18. Temperature dependence of Y123 (upper Ag-doped Y123 (lower curve) films at zero field.

curve) and

a sample with smaller LY corresponds to smaller 7. These results suggest that the samples with better pinning have smaller (Y(and/or 7). We believe that the difference in the detailed microstructure is the key factor in this observation. A more detailed comparison of the Y123 and Agdoped Y123 superconducting films revealed that c-axisoriented films prepared from the Ag-added Y123 target exhibit better superconducting characteristics, including transition temperature and critical current densities. TEM studies show that the interface between the Agdoped Y123 film and the (001) orientation MgO substrate is much sharper and exhibits a periodic interfacial dislocation. Figure 18 gives the typical temperature dependence of resistivity curves for both the Y123 and Ag-doped Y123 films. The Y123 films have resistivities

close to 4~ lop4 R-cm at 200 K, which is about a factor of 5 larger than that of single crystals. The resistivity for the Ag-doped Y123 film is about 60% that of the Y123 film. SEM pictures of the films show no obvious visual differences in microstructure between the Y123 and Ag-doped Y123 films. X-ray diffraction patterns of the Y123 film and the Ag-doped Y123 film show no impurity phases (except Ag metal). These observations suggest that silver does not go into the Y123 lattice, which is consistent with what was observed in the bulk AgN123 composite. Figure 17 clearly shows that the Ag-doped film has a larger slope, in comparison with other films of similar thickness, which suggests that its pinning strength is much larger. It is expected that there is an enhancement in the critical current densities for the Ag-doped film. Indeed, the J, values, determined either from the direct transport or from magnetic hysteresis measurements at 77 K in zero field, for the Ag-doped Y123 films are about an order of magnitude larger than those of the undoped samples. High-resolution TEM images of the Y123 and Ag-doped Y123 films are shown in Fig. 19. The pictures clearly show that the interfaces between the film and the substrate are quite different for the two films. The interface for the Ag-doped film (Fig. 19(a)) is very sharp and shows interfacial dislocation with = 6 A periodicity. The Burgers vector of these dislocations is determined to be l/2 [OlO],,,. On the other hand, the interface for the undoped Y123 (Fig. 19(b)) shows amorphous-like structure. It is not clear at present what the role of the Ag is in leading to the above observation. One plausible source is likely the lattice mismatch between Ag and the MgO substrate. This observation is believed to have direct consequences for the enhancement of the observed properties. The existence of a true superconducting state (or vortex glass state) in the Y123 system is clearly demonstrated from our observations. The enhancement in critical current densities and the materials formability can then be achieved by properly controlling the materials processing conditions and the introduction of dopants such that the effect to the intrinsic superconducting characteristics remains minimal. It is obvious that a more detailed theoretical picture is required in order to fully understand the above observations. More complete work using a wider selection of materials and better processing techniques is also needed in order to achieve practical applications of the high T, superconductors. The mystery of the resistive ‘knee’ observed in the crystal and thick film samples remains unexplained. However, preliminary results or our recent study suggest that this unusual behavior in the Y123 system may correlate with the dimensional crossover of the vortex motion in the liquid state. This anomaly may also strongly relate to the observed difference in

198

large-scale applications more viable, Information needed to understand the vortex dynamics in the oxide at low temperature and high magnetic field has been accumulated. Experimental data and the theoretical picture are starting to merge, though a complete understanding of the phenomena requires more detailed, more accurate, better-characterized materials and more quantitative theory. Decoding the mystery of the superconducting mechanism requires a more decisive experiment. Although many important problems remain unresolved, the progress in the last five years is unprecedented, and it is expected that the same level will be maintained in the future.

Acknowledgements The author thanks Professor J. T. Chen, Dr C. C. Chi, Dr Y. Huang, Dr H. Y. Tang, Dr S. R. Sheen, Dr C. H. Kao, Prof. F. Z. Chicn, Mr M. J. Wang, Mr C. L. Lin, Mr C. C. Chen and other members of the Superconductivity Laboratory in the Materials Science Center at the Tsing Hua University for their valuable contributions. He also thanks Dr C. H. Chen, Dr L. J. Chen and Dr F. R. Chen for their TEM work. The work was supported by the RUC National Science Council, grant NSC81-0511-M007-02.

Fig. 19. TEM images (cross section} of (a} Ag-doped (b) Y123 thin films on (001) MgO substrate.

J, from magnetic surement.

measurement

Y123 and

and transport

mea-

6. Conclusions

Problems related to the phase stabilization, including oxygen stabil~ation~ of the high-temperature superconducting oxides are discussed in this article. I have also presented the possibility of using novel processes to achieve low-temperature oxidation in these materials. These developments open up not only the possibility of more controllable materials processing techniques, but also channels for the search for new high T, materials. Novel approaches using either metal-superconductor composites or the electrochemical crystallization technique have successfuily been applied to grow large, high-quality superconducting single crystals. These provide great promise for future direction in terms of continuous texture or low-temperature processing possibilities, which in turn make the future

J. G. Bednorz and K. A. Muller, 2. P&s. l3, 64 (1986) 189. M. K. Wu, J. R. Ashburn, C. J. Tomg, P. H. Hor, R. L. Meng, L. Gao, Z. J. Huang, Y. Q. Wang and C, W. Cbu, Plsys. Rev. Len, 58 (1987) 908. See, for example, H. Iiukuyama, S. Maekawa and A. P. Malozemoff (eds.), SfruBg Correta&‘om and S~pe~cu~d~c?~~~~, Springer Series in Solid State Science, Vol. 9, Springer, Heidelberg, 1989; S. Maekawa and M. Sato (eds.), Physics of High Temperature Superconductws, Springer Series in Solid State Science, Vol. 106, Springer, Heidelberg, 1991; D. M. Ginsberg (ed.), Physical Propetiies of High Temperature Superconductors, Vols. 1 and 2, World Scientific, Singapore, 1990 and 1991. J. Bardeen, L. N. Cooper and 3. R. Schrieffer, Phys. Rev., 108 (1957) 1175. H. Maeda, Y. Tanaka, M. Fukutomi and T. Asano, Jpn. J. Appi. Phys., 27 (1988) L.209. Z. Z. Sheng, A. M. Herman, A. El Ali, C. Almason, J. Estrada, T* Datta and R, J, Matson, P&s, Reu, Lett., 50 (1988) 937” R. J. Cava, B. Batlogg, J. J. Krajewski, R. Farrow, L. W. Rupp, Jr., A. E. White, K. Short, W. F. Peck and T. Kometani, Nature (London), 332 (1988) 814. B. BatIogg, R. J. Cava, A. Jayarama, R. B. van Dover, G. A. KourouMis, S. Sunshine, D. W. Murphy, L. W. Rupp, Jr., H. S. Chen, A. White, K. T. Short, A. M. Mujsca and E. A. Rietman, Whys. Rev. ten., 58 (1987) 2333; L. C. Boume et al_, Whys Rev. Lert., 58 (1987) 2337; B, Batlogg, R. J. Cava,

199 L. W. Rupp, Jr., A. M. Mujsce, J. J. Krajewski, J. P. Remeika, W. F. Peck, Jr., A. S. Cooper and G. P. Espinosa, Phys. Rev.

37 K. Fueki,

L&t.,

38 E. J. M. O’Sullivan and B. P. Chang, Appl.

61 (1988)

1670.

9 J. R. Schrieffer, S. G. Wen and X. C. Zhang, Phys. Rev. Lett., 60 (1988) 944. 10 P. W. Anderson, Science, 235 (1987) 1196. 11 V. J. Emery, Phys. Rev. Lett., 58 (1987) 2794. 12 J. E. Hirsch, Phys. Rev. Lett., 59 (1987) 228. 13 G. Chen and W. A. Goddard III, Science, 239 (1988) 899. 14 R. Friedberg and T. D. Lee, Mod. Phys. Lett., B2 (1988) 469; T. D. Lee, Proc. Con& Two-Dimensional Strongly Correlated Electronic Systems, Beijing, May, 1988, p. 113; T. D. Lee, Nature (London), 330 (1987) 460. 15 C. C. Tsuei, D. M. Newns, C. C. Chi and P. C. Pattnaik, Phys. Rev. Lett., 65 (1990) 2724; C. C. Tsuei, D. M. Newns,

C. C. Chi, P. C. Pattnaik and M. Daumling, preprint, 1992. 16 M. Garvith and A. T. Fiory, Phys. Rev. Lett., 59 (1987) 1337. 17 T. R. Chien, Z. Z. Wang and N. P. Ong, Phys. Rev. Len., 67 (1991)

2088.

18 H. Ishi, H. Sato, N. Takagi, S. Uchida, K. Kitazawa, K. Kishio, K. Fueki and S. Tanaka, Physica B, 148 (1987) 419. 19 A. R. Moodenbaugh, Y. Xu, M. Suenaga, T. J. Folkerts and R. N. Shelton, Phys. Rev. B, 38 (1988) 4596. 20 J. D. Axe, D. E. Cox, K. Mohanty, H. Moudden, A. R. Moodenbaugh, Y. Xu and T. Thurston, IBM J. Res. Dev., 33 (1989)

382.

21 Y. Uemura, L. P. Le, G. M. Luke, B. J. Stemlieb, W. D. Wu, J. H. Brewer, T. M. Riseman, C. L. Seaman, M. B. Maple, M. Ishikawa, D. G. Hinks, J. D. Jorgensen, G. Saito and H. Yamaochi, Phys. Rev. Lett., 66 (1991) 2665. 22 N. P. Ong, in D. M. Ginsberg (ed.), Physical Properties of High Temperature Superconductors, Vol. 2, World Scientific, Singapore, 1991, p. 459. 23 P. Chaudhari, R. H. Koch, R. B. Laibowitz, T. R. McGuire and R. J. Gambino, Phys. Rev. Lett., 58 (1987) 2684. 24 M. K. Wu and J. R. Ashburn, unpublished; R. Pool, Science, 241 (1988)

655.

25 M. K. Wu, J. R. Ashburn, C. A. Higgens, C. W. Fellows and C. Y. Huang, Appl. Phys. Lea, 52 (1988) 1915. 26 M. K. Wu, J. R. Ashburn, C. A. Higgins, B. H. Loo, D. H. Burns, A. Ibrahim, T. Rolin and C. Y. Huang, Phys. Rev. B, 37 (1988)

3678.

30 J.-C. Grenier, A. Wattiaux, N. Laguyete, J. C. Park, E. Marquestaut, J. Etourneau and M. Pouchard, Physica C, I73 (1991) 139. 31 H. Y. Tang, W. L. Chen and M. K. Wu, to be published. 32 See, for example, M. G. Smith, A. Manthiram, J. Zhou, J. B. Goodenough and J. T. Markert, Nature (London), 351 (1991) 549. 33 M. Takano, M. Azuma, Z. Hiroi, Y. Bando and Y. Takeda, Physica

C, 176 (1991)

441.

34 J. T. Chen, L. X. Qian, L. Q. Wang, L. E. Wenger and E. M. Logothetis, Mod. Phys. Lett., B3 (1989) 178. 35 M. K. Wu, D. C. Ling, M. J. Wang and J. T. Chen, Proc. NASA Conj: High MD, 1990, p. 1.

Temperature

Superconductivity,

Greenbelt,

36 K. N. Tu, C. C. Tsuei, S. I. Park and A. Levi, Phys. Rev. B, 38 (1988)

772.

K. Kitazawa,

State

(1988)

K. Kishio and T. Hasegawa,

Sci., 2 (1988)

Rev.

219. Phys. Lett.,

52

1441.

39 M. K. Wu and D. C. Ling, Chin. J. Phys. (Taipei),

28 (1990)

349.

40 M. Murakami, S. Gotoh, N. Koshizuka, S. Tanaka, T. Matsushita, S. Kambe and K. Kitazawa, Cryogenics, 30 (1990) 390. 41 D. L. Kaiser, F. Holtzberg, M. F. Chisholm and T. K. Worthington, J. Cryst. Growth, 85 (1987) 593. 42 M. K. Wu, M. J. Wang, J. L. Lin, D. C. Ling and K. J. Ma, AIP

Confi Proc.,

219 (1991)

100.

43 Y. Huang, Y. F. Lin, M. J. Wang, M. K. Wu and K. J. Ma, preprint. 44 M. L. Norton and H. Y. Tang, Chem. Mater., 3 (1991) 431. 45 H. Y. Tang, W. L. Chen and M. K. Wu, preprint; S. F. Lee, J. Y. T. Wei, H. Y. Tang, T. R. Chien and M. K. Wu, preprint. 46 D. G. Hinks, B. Dabrowski, J. D. Jorgensen, A. W. Mitchell, D. R. Richards, Shiyou Pei and Donglu Shi, Nature (London), 333 (1988)

836.

47 P. H. Hor, R. L. Meng, Y. Q. Wang, L. Gao, Z. J. Huang, J. Bechtold, K. Forster and C. W. Chu, Phys. Rev. Left., 58 (1987)

1981.

48 See, for example, J. T. Markewrt, Y. Dalichaouch and M. B. Maple, in D. M. Ginsberg (ed.), Physical Propetiies of High Temperature Superconductors, Vol. 1, World Scientific, Singapore, 1990, p. 293. 49 B. W. Veal, W. K. Kowk, A. Umezawa, G. W. Crabtree, J. D. Jorgensen, J. W. Downey, L. J. Nowicki, A. W. Mitchell, A. P. Paulikas and C. H. Sowers, Appl. Phys. Lett., 51 (1987) 279.

50 L. Soderholm, K. Zhang, D. G. Hinks, M. A. Beno, J. D. Jorgenson, C. U. Segre and I. K. Schuller, Nature (London), 328 (1987)

604.

51 J. Wei, W. C. Chen, W. Y. Guan and M. K. Wu, unpublished results. 52 M. Oda, T. Murakami, Y. Enomoto and M. Suzuki, Jpn. J. Appl.

Phys.,

26 (1987)

L804.

53 Q. R. Zhang, L. Z. Cao, Y. T. Qian, Z. Y. Chen, W. Y. Guan, Y. Zhao, G. G. Pang, H. Zang, J. S. Xia, M. J. Zhang, D. Q. Yu, Z. H. He and S. F. Sun, Solid State Commun., 63 (1987)

9765.

27 Yunhui Xu and W. Y. Guan, Physica C, 183 (1991) 105. 28 B. Raveau, C. Michel, M. Hervieu, J. Provost and F. Studer in J. G. Bednorz and K. A. Muller (eds.), Earlier and Recent Aspects of Superconductivity, Springer Series in Solid State Science, 90, Springer, Heidelberg, 1990, p. 66. 29 J. M. Sanchez, F. Meja-Lira and J. L. Moran-Lopez, Phys. Rev. B, 37 (19880

Solid

535.

54 B. Jayaram, Solid State

S. K. Agarwal, Commun.,

A. Gupta

63 (1987)

and A. V. Narlikar,

713.

55 R. J. Cava, B. Batlogg, J. J. Krajewski, L. W. Rupp, L. F. Schneemeyer, T. Siegrist, R. B. van Dover, P. Marsh, W. F. Peck, Jr., P. K. Gallagher, S. H. Glarum, J. H. Marshall, R. C. Farrow, J. V. Waszczak, R. Hull and P. Trevor, Nature (London), 336 (1988) 211. 56 S. Adachi, K. Setsune and K. Wasa, Jpn. J. Appl. Phys., 29 (1990)

L890.

57 S. F. Hu, D. A. Jefferson, R. S. Liu and P. P. Edwards, preprint. 58 M. K. Wu, D. C. Ling, M. J. Wang and F. 2. Chien, in H. S. Kwok and D. T. Shaw (eds.), Superconductivity and Its Applications, Plenum, New York, 1990. 59 S. A. Sunshine, L. F. Schneemeyer, T. Siegrist, D. C. Douglass, J. V. Waszczak, R. J. Cava, E. M. Gyorgy and D. W. Murphy, Chem.

Mater.,

I (1989)

331.

60 J. L. Hsu. Master’s Thesis. Tsine Hua Universitv. 1991. 61 N. Mori, H. Takahashi, Y. Shimakawa, T. Ma&k0 and Y. Kubo, J. Phys. Sot. Jpn., 59 (1990) 3839. 62 H. Yamamoto, T. Mori, S. Onari and T. Arai, Physica C, 181 (1991)

133.

63 Y. Matsuda, M. Yoshida and S. Hinotani, Jpn. J. Appl Phys., 28 (1991) L1128. 64 M. Hangyo, S. Nakashima, M. Nishiuchi, K. Nii and A. Mitsuishi, Solid State Commun., 67 (1988) 1171.

200 65 Y. Morioka, M. Kikuchi and Y. Syono, Jpn. J. Appt. Phys., 26 (1991) L1499. 66 M. Cardona, L. Genzel, R. Liu, A. Wittlin, H. J. Mattausch, F. Garcia-Alvarado and E. Garcia-Gonzalez, Solid State Commun., 64 (1987) 727. 67 S. F. Lee, J. Wei, D. H. Chen, S. R. Sheen and M. K. Wu, to be published. 68 D. Esteve, J. M. Martinis, G. Urbina, M. H. Devorat, G. Collin, P. Monod, M. Ribault and A. Revcolevschii, Europhys. Lett., 3 (1987) 1237. 69 G. Deutscher and K. A. Muller, Phys. Rev. Lett., 59 (1987) 1745. 70 B. Roas, L. Schultz and G. Saemann-Ischenko, Phys. Rev. Lett., 64 (1990) 479. 71 P. N. Peters, R. C. Sisk, E. W. Urban, C. Y. Huang and M. K. Wu, Appl. Phys. Lett., 52 (1988) 2066. 72 C. Y. Huang, H. H. Tai and M. K. Wu, Mod. Phys. Lett., B3 (1989) 525. 73 J. K. Tien, J. C. Borofka, B. C. Hendrix, T. Caulfield and S. H. Reichman, Metall. Trans. A, 19 (1988) 1841. 74 B. C. Hendrix, J. C. Borofka, T. Abe, J. K. Tien, T. Caulfield and S. H. Reichman, in W. Mayo (ed.), Processing and Applications of High T, Superconductors: Status and Prospects, Proc. 1988 TMS Northeast Regional Meet., Rutgers University, May 9-11, 1988, p. 169. 75 B. C. Hendrix, J. C. Borofka, T. Abe, J. K. Tien, T. Caulfield and S. H. Reichman, in H. S. Kwok and D. T. Shaw (eds.), Superconductivity and Its Applications, Elsevier, New York, 1988.

76 K. A. Muller, M. Takashige and J. G. Bednorz, Phys. Rev. Lett., 58 (1987) 1143. 77 Y. Yeshurun and A. P. Malozemoff, Phys. Rev. Lett., 60 (1988) 2202. 78 Y. Iye, T. Tamegai, T. Sakakibara, T. Goto, N. Miura, H. Takeya and H. Takei, Physica C, 153-155 (1988) 26. 79 T. T. M. Palstra, B. Batlogg, R. B. van Dover, L. F. Schneemeyer and J. V. Waszczak, Appl. Phys. Lett., 54 (1989) 763. 80 P. L. Ganunel, L. F. Schneemeyer, J. V. Waszczak and D. J. Bishop, Phys. Rev. Lett., 61 (1988) 1666. 81 R. H. Koch, V. Foglietti and M. P. A. Fisher, Phys. Rev. Lett., 64 (1990) 2586. 82 P. L. Gammel, L. F. Schneemeyer and D. J. Bishop, Phys. Rev. Lett., 66 (1991) 953. 83 T. K. Worthington, E. Olsson, C. S. Nichols, T. M. Shaw and D. R. Clarke, Phys. Rev. B, 43 (1991) 10538. 84 M. Tinkham, Phys. Rev. Lett., 61 (1988) 1658. 85 D. R. Nelson and H. S. Seung, Phys. Rev. B, 39 (1989) 9153. 86 D. S. Fisher, M. P. A. Fisher and D. A. Huse, Phys. Rev. B, 43 (1991) 130. 87 S. P. Obukhov and M. Rubinstein, Phys. Rev. Lett., 65 (1990) 1279. 88 M. V. Feig’lman, V. B. Geshkenbein, A. I. Larkin and V. M. Vinokur, P&s. Rev. Lett., 62 (1989) 2857. mm~mm89 M. K. Wu, W. J. Wang, C. C. Chi and F. Holtzberg, Chin. J. Phys. (Taipei), 30 (1992) 253. 90 M. K. Wu, W. J. Wang, T. R. Chien, Y. C. Chen, C. C. Chi and F. Holtzberg, in Proc. Mater. Res. Sot. 1992 Spring Meet., San Francisco, Apr. 1992, p. 201.