Superconducting properties of A15 V3Ga and Nb3Al by precipitation from supersaturated bcc phase

Structu rally multiconnected filamentary A 15 V3Ga and Nb3AI in their respective bcc phases are formed by heat treating as quenched and deformed supersaturated bcc phases. It was found that Tc of these compounds are " 1 4 and ~12 K V3Ga and Nb3AI, respectively, after appropriate heat treatments. The kinetics and morphology of precipitation of these compounds as well as the critical current density are presented.

Superconducting properties of A15 V3Ga and Nb3AI by precipitation from supersaturated bcc phase V.M. Pan In recent years, a number of large scale uses of superconductors for technological projects have been planned and constructed. Requirements for superconductors in these systems are at times conflicting. For example, for increased stability, a larger portion of the conductor has to be devoted to a stabilizer which for a higher field operation, a high overall current density is needed. A possible solution to this problem is based on drawing the superconducting fibres to their smallest possible diameter; in this way, it is possible to satisfy the above requirements to a certain degree. If the filling factor remains constant, the overall current density must be increased because of the increased surface area of the superconductor. This also results in a greater degree of stabilization and lower losses in alternating magnetic fields. Under these circumstances the superconductor becomes more susceptible to bending deformations with no degradation of the critical current. The technology to produce traditional A15 multifllamentary composite superconducting materials was first proposed in 1971.1'2 Useful superconducting filaments as small as 5/am in diameter were obtained. Smaller filament diameters resulted in much more frequent fracture of filaments, which had a strongly adverse effect on the composite superconducting properties. Naturally, physicists and technicians are lookhag for ways to fabricate superconducting materials by other methods. In particular, they are searching for methods of creating the discrete structure of superconducting current paths within the normal matrix by the use of the natural physical-chemical processes rather than by the 'forced' means used for producing traditional multifilamentary composite superconductors. At present, there are several methods of producing A15 superconducting materials with a discrete structure of superconducting current paths ('channels'). These materials are now known as 'in-situ' prepared superconductors, but this term is hardly acceptable. The label 'structurally multiply connected superconductors' appears to be more proper and will survive longer.

Structurally multiply connected superconductors All the presently known versions of materials with a multiple connected discrete structure of superconducting channels in a normal (or 'poorly"superconducting) matrix can be subdivided into three groups, according to the method of processing and into two types, according to the degree of The author is from the Institute of Metal Physics, 36 Vernadsky Str., 252680, Kiev 142, USSR. Paper received 27 February 1981.

structural dispersion with all the subsequent physical consequences. The first group is distinguished by the method of crystallization: either directional 3 or random 4~5 crystallization of two- or three-component alloys (Cu-Nb, Cu-V, Cu-Nb-Sn, Cu-V-Ga) is used. Alloys are subjected to deformation through rolling or drawing after casting and diffusion annealing to get the directional multiply connected structure of A15 phases. The second group of materials uses powders of two or three components (eg, Nb, Cu, Sn). The powders are subjected to hot 7 or cotd 8 extrusion followed by diffusion annealing to produce the directional multiple connected structure of A15 phases. Preliminary inf'dtration of liquid Sn is also used in porous powder rods of Nb with subsequent deformation and annealing of the composite (the so-called 'infdtration method'). 9 tn the third technique, the main fabrication process, is the aging of a supersaturated solid solution to precipitate a new phase (A15) in a solid state, l°'la The first two groups differ by the degree of dispersion in the system of multiple connected superconducting channels. The typical cross sectional size of formations of the supercon ducting A15 phase is in the range 103 to 104 ,8, For materials prepared by the powder metallurgy method 7,s the filament size is larger, for those obtained by the method of casting and subsequent deformation the filament size is smaller (sometimes as small as 200 A).6 However the characteristic filament size, d, of the superconducting A15 phase can never be as small as the coherency length, ~ (d > ~). In contrast, for materials of the third group prepared by the aging of a supersaturated solid solution (described below), the typical size of a superconducting filament is about 102 A and less; therefore, the condition d ~ ~ can be satisfied for these materials. As is known, if d ¢ ~, the destruction of superconductivity by a current in the presence of magnetic field occurs because the superconducting condensate has reached its critical velocity, not because of vortex instability. This was shown by Bardeen as early as 1962.14 The depairing criterion connected with the current at T = 0 determines the critical current density, Jc, as follows: Je "Z'_0.82 Hc (1) 4trekL where Hc is the thermodynamical critical magnetic field, and XL is the London's penetration depth. In this case Jc

0011-2275/81/009547-08 $02.00 © 1981 I PC Business Press Ltd CRYOGENICS. SEPTEMBER 1981

547

was estimated as 100 MA cm -2. The depairing mechanism is realized when the vortex instability is suppressed for some reason, eg, when vortices do not fit into the superconductor because of its small size (d ~ ~). The method of solid state precipitation for production of multiple connected superconductors has the advantage of being able to produce the narrowest superconducting channels. To actually produce such materials, the following conditions must be satisfied: one, the A15 phase must be in equilibrium with the bcc solid solution (on the base, eg, of Nb or V) or with the fcc solid solution (on a Cu base); two, solubility of the B component (B = A1, Ga, Si, Ge, Sn) in vanadium, niobium, or copper must be strongly temperature dependent, otherwise the required supersaturation of a metastable solid solution becomes impossible; three, the supersaturated metastable solid solution must not decompose when supercooled; four, the solid solution supersaturated with the B component must remain sufficiently ductile; five, the intermetallic A15 phase that is formed during aging must attain a stoichiometric composition; and six, the optimum structure, in terms of the amount, size and arrangement of the precipitating A15 phase, must be created during the aging of the supersaturated solid solution.

Properties of V-Ga and Nb-AI deformed aging alloys As early as 1969, we performed the first experiments on the preparation of structurally multiple connected superconductors using the aging process of a deformed supersaturated solid solution of a bcc lattice in V-Ga alloys, l° Later we obtained similar superconductors with Nb-A1 alloys. Then we studied the structure and superconducting properties of both of these alloy systems. The process developed to produce narrow superconducting 'channels' of the A15 phase consists of melting, fast cooling, deformation, and thermal treatment. The state of a homogeneous solid solution with a bcc phase on a vanadium or niobium base is determined by the initial composition of the alloy (up to 16-17 arm% Ga or 13-14 arm% A1). The V-Ga alloys were also alloyed with cerium (about 0.8 at m%) to improve the alloy ductility with interstitial impurity bonding and to ensure the stoichiometry of the precipitating VaGa phase according to the phase equilibrium diagram for the V-Ga-Ce system.15 In this initial state, the alloy is subjected to plastic deformation with a high degree of compression (e ~ 8090%). The dislocation structure formed as a result has a high dislocation density (up to 1011-1012 cm-2). The details of the structural and superconducting properties of these materials is described in the following sections.

Dislocation structure peculiarities. Note that in the early stages of V-Ga alloy deformation, there is a relatively homogeneous dislocation distribution) 6 Dislocations are aligned mainly along a screw component. The observed shape of the dislocation lines indicates a relatively low stacking-fault energy, with that of metals belonging the VIA group for example. This is due to the slower propagation of screw dislocations. Dislocation density grows steadily with the increase in deformation, and single dislocations become impossible to resolve. At large deformations (e 2 90%), the structural state appears at which plastic deformation occurs exclusively on highly curved sliding planes. This is proved by a strong azimuthal blurring of reflexes on electron pictures. Such structural states have a special feature - improvement of strength characteristics along with a decrease in the trans-

548

ition temperature to brittle fracture. It must be said that inhomogeneous dislocation distribution is observed in some parts of the plane normal to the rolling direction. This suggests the onset of a cellular structure aligned along the rolling direction. Deformation texture seems to be the reason for such inhomogeneous distribution. At high deformations (e ~ 80%), the Nb-A1 alloys have a distribution structure in the rolling plane that is less homogeneous than that in V-Ga alloys. Oblong free cells, 0.2 to 0.3/am in size, are formed during deformation in the normal plane.

Structure and phase changes at annealing. Subsequent thermal treatment of alloys (aging annealing) leads to decompo. sition (aging) of a deformed, supersaturated, metastable solid solution and the precipitation of A15 phase disperse formations, which decorate the alloy dislocation structure. Decoration occurs in such a way that a multiple connected three-dimensional network of mainly continuous, narrow, superconducting channels is created in a normal (more exactly, 'poorly' superconducting) matrix. When deformed alloys are annealed, the dislocation structure is reconstructed. In V-Ga alloys this structure is already rather unstable during the process of polygonization at short (15-30 rain) annealings at 650 to 800°C. The texture developed during the rolling is very important during polygonization. If the rolling plane contains planes of the I 1111 type, a perfect equiaxial cellular structure is formed, but when it contains the planes of the [ 100 1 type, either oblong, weakly misoriented sub-boundaries and dislocation networks are formed or the inhomogeneous dislocation distribution appears in elongated parts 10/.tin in size. In Nb-AI alloys, the increase in aluminium concentration strongly hampers the dislocation structure evolution (cellular structure formation, polygonization, and recrystallization). Annealing at 950°C for 1 h has practically no effect on the character of the dislocation structure in the alloy Nb-13 atm% A1. Dislocation redistribution and formation of cells in the c r o s s section parallel to the rolling plane are not observed. The process of cell formation in this alloy begins only after 5 h of annealing at 950°C; the well-developed equiaxial cellular structure appears after 10 h of annealing. Comparison of data obtained by the TEM method, by measuring of electrical resistance and superconducting properties, suggests that aging annealing does not lead to the coherent homogeneous nucleation in V-Ga and Nb-A1 alloys. The A15 phase is precipitated by a heterogeneous mechanism through nucleation on defects, dislocation walls and sub-boundaries, grain boundaries, and s o o n . 16'17

The phenomenon of precipitate "invisibility'. There is a common and rather interesting feature - the 'invisibility' of VaGa and Nba A1 phase precipitates when they are studied with an electron microscope in the early stages of aging, although the superconducting properties of aged alloys strongly suggest the presence of the A15 phase. The most interesting and demonstrative property effect is critical current variation with time (Fig. 1) in V-Ga samples of different compositions deformed at 4.2 K and subjected to isothermal annealing at 800°C. The matrix (a solid solution with a bcc phase) is not superconducting at 4.2 K. Fig. 2 shows the dependence of the critical temperature of a bcc solid solution on gallium content. The V3 Ga precipitation is undoubtedly responsible for the appearance of superconductivity (after annealing), with a 14 K critical

C R Y O G E N I C S . SEPTEMBER 1981

microscope) have revealed any traces of phase precipitation at this stage. Invisibility may perhaps be explained by the small characteristic size of precipitates (less than 70-100 ~k). Precipitate presence on dislocations is masked by the contrast of the dislocation line itself. From elastic stress fields near the line, its width is observed to be about 100 -A. Furthermore, the precipitates on dislocations are not coherently connected with the matrix and do not create elastic deformations, which might have enhanced the contrast and thus increased the probability of observing the phase.

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A similar situation exists in Nb-A1 alloys, with the exception of Nb-A1 solid solutions with a bcc lattice (T O = 5.5 K) where at 4.2 K the Nb3A1 phase precipitates are not in the normal but in the superconducting matrix. Therefore, in this case it is impossible to follow clearly the appearance of superconductivity during aging annealing, as it was in Fig. 1, and despite high critical current densities in magnetic fields (at least up to 6 T), it has been impossible to discover any traces of a precipitated phase, even with the TEM method.

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61

Precipitation of A15 phases may be observed only during over-aging, when these phases become much larger. Coarse formations (d _~ 104 A) of the V3Ga phase precipitates are observed in alloys with 16 atm% Ga after long annealing at 800°C. Since these particles are separated from one another by macroscopic spaces, they do not significantly contribute to the critical current magnitude. When the same sample is looked at in a dark field using one of the V3 Ga diffraction spots, very thin layers of precipitate could be seen on the boundary of a dislocation cell. Hence, it may be assumed that these alloys can carry superconducting current because during annealing a multiple connected threedimensional network is formed composed of A15 phase continuous particles with a characteristic cross section equal to or less than the coherency length (d ~ ~). Besides the phenomenon of the 'invisibility' of A15 phase precipitates, other peculiarities in properties have also been discovered in both the systems that occur during aging annealing of deformed V-Ga and Nb-A1 alloys.

V-Ga alloys. During early stages of the aging annealing, an increase in the electrical resistance, AR/Ro, is observed in the normal state. It depends on the gallium concentration in the alloy (the more Ga, the higher AR/Ro) and the degree of deformation before annealing. This effect is practically absent in recrystallized alloys, but amounts to 5-7% where there is a high degree of deformation (Fig. 3). Perhaps during the early stages, gallium atoms segregate (deposit) on dislocations, which results in formation of Cottrell's 'Clouds' or 'Pipes'.

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temperature and densities of the critical current above 100 kA cm -2 . But careful investigations of the aged alloy structure by TEM methods (including those in a dark field as well as by the method of a weak beam in a million-volt

CRYOGENICS. SEPTEMBER 1981

The critical temperature, Tc, of the V3 Ga phase created during aging depends strongly on the degree of deformation of the initial bcc solid solution. The T c varies from 5 to 6 K for cast and weakly deformed alloys and to 14 to 15 K for 70 to 80% reduction in area (Figs 4 and 5). This phenomenon, along with the increase in the electrical resistance during the early aging stages, is apparently connected with the segregation of gallium atoms on dislocations or cell boundaries. The maximum on the Tc-e curve for V-Ga alloys (Fig. 5) means that, owing to segregation, the local gallium concentration grows with the increase in deformation and approaches stoichiometric composition, which ensures a higher T c. But at very large deformations, gallium segregation attains such a level that the stoichiometric ratio between the com-

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ponents in the precipitating V3 Ga phase is exceeded, and hence Tc drops. Since the precipitating A15 phase nucleates on dislocations, the process of thermally activated dislocation motion during annealing must influence the kinetics and morphology of the precipitates. Two mechanisms are possible: the climbing dislocations are released from the cloud of impurity atoms and the new-phase particles sitting on it or the climbing dislocations drag along the particles of the precipitated phase. From our data on structural investigations and on electrophysical and superconducting properties of deformed alloys i_n the annealing process, we assume that both of these mechanisms participate to some extent. The nuclea-

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tion process involved when the new-phase nucleus trails behind the moving dislocation may be represented as follows: precipitation nuclei on dislocations are very small; where there are no dislocations, the existence of a'nucleus of such size is unlikely because of the energy requirements; therefore when dislocations are moving, the precipitates at the rear dissolve while precipitates at the front grow. In the process of polygonization, the moving dislocation can 'carry over' the new-phase nuclei onto cell boundaries. When climbing dislocations break away from sessile particles of a precipitating phase (or annihilate them), the precipitation dissolves. The redissolution of some of the V3 Ga phase is perhaps responsible for the decrease in the critical current after annealing at 800°C for 20 to 30 minutes. The presence of two regions of higher T c separated by a marked minimum strongly suggests an alteration of the precipitate character. The initial increase in the critical current (Ic) is apparently due to a large number of nuclei of a new precipitating phase. These nuclei appear on dislocations and do not appreciably grow. But this process is exhausted when the dislocation concentration decreases and a somewhat stable cellular substructure forms. The second region of higher T c results from the growth of the precipitates on dislocation sub-boundaries. The minimums on the lc (r) curves may be attributed to the partial dissolution of those precipitates of less-than-critical size (they did not enter the subboundaries and were left in a cell volume after the dislocations had gone away or were annihilated). In deformed alloys, the process of V3 Ga precipitation is much slower than that of a dislocation structure reconstruction and polygonization at the same annealing temperature. Therefore the critical current density depends rather weakly on the degree of deformation.

Nb-AI alloys. The growth of electrical resistance in a normal state during the early stages of deformed alloy annealing has not been observed (Fig. 6). The critical temperature (11 to 12 K) of the Nb3A1 phase formed during aging is nearly

CRYOGENICS

. S E P T E M B E R 1981

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independent of the degree of deformation of the initial solid solution with a bcc lattice (Fig. 5). Such a Tc value corresponds to the Tc for a steady state nonstoichiometric A15 phase enriched with niobium and found in the twophase structure (bcc solution plus A15 phase-Nb0.Ts+x Alo.2s -x) at the annealing temperature (950°C) and is much below the T c value for the stoichiometric Nb3 A1 (18.5 K). Apparently, segregation ('deposition') of aluminium atoms on dislocations or dislocation formations does not take place. However, it is quite possible that niobium atoms segregate on defects, and that the resulting increase in the electrical resistance is masked by its fast decrease due to precipitation of the second phase. In deformed alloys, the formation of the Nb3A1 is much faster than the development and polygonization of the dislocation structure at the same annealing temperature. Therefore, the critical current density grows sharply as the deformations increase (Fig. 7). Critical current. An important characteristic of structurally

multiple connected superconductors (of other superconducting materials also) is the 'riffling factor', that is, the ratio of areas filled by the superconductor and the normal metal in the sample cross section. The larger the filling (or packing) factor the higher, naturally, is the overall critical current (other conditions being equal) in the material. It is clear that the filling factor should be increased to produce the largest possible number of ultimately narrow superconducting 'channels'. Therefore, for materials of this type the following condition is of extreme importance: the stability of the dislocation structure during reconstruction and polygonization must be much greater than that of the supersaturated solid solution at the same annealing temperature. As has

CRYOGENICS.

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been shown above, this condition applied to Nb-A1 alloys but not to V-Ga alloys, where the critical current density is lower for just this reason. Fig. 8 shows the V-Ga critical current as a function of a magnetic field, H. Note the small difference between the Ic (H) values for the magnetic field parallel and perpendicular to the current. This may be due to the destruction of superconductivity which occurs, not because of vortex instability, but because the superconducting condensate has reached its critical velocity (depairing) Fig. 9 gives critical current density as a function of the field for the three samples of Nb-A1 alloys with the same degree of deformation (85-88%) aged at 950°C for different time periods. Curve 1 refers to the moment when the 1- V characteristics begin to depart from the x axis on the recorder; curve 3 refers to thermal breaking away in a normal state. The critical current density amounts to 25 kA cm -2 in the perpendicular magnetic field of 6 T at 4.2 K. is In the same figure, the 1-V characteristics are shown for these samples when the superconductivity is destroyed by the current. The existence of a somewhat significant resistive region between zero and normal resistances was quite unexpected. This means that there exists some mechanism of energy dissipation that does not cause the transition into a normal state. The presence of a resistive state in the alloy with narrow superconducting channels may have two possible explanations. The first one is trivial and consists of the fact that channels contain inhomogeneities in their cross sections and the con1000

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Fig. 9 a -- Critical current density at 4.2 K vs the magnitude of the external perpendicular magnetic field. I-V characteristics of s u p e r c o n d u c t i v i t y destruction b y the current of the alloy Nb-13 atm % A 1 . a - d e f o r m a t i o n b y 85%, annealing at 9 5 0 ° C f o r 3 h; b -- d e f o r m a t i o n b y 85%, annealing at 9 5 0 ° C for 1 h: c - - d e f o r m a t i o n b y 50%, annealing at 1800°C f o r 5 rain. fast cooling, d e f o r m a t i o n by 88%, annealing at 9 5 0 ° C for I h. 1' -- the beginning o f departure of the I-V curve f r o m the x axis; 2" -- the appearance of a potential difference of 0.1 mW on the sample; 3' -- the rush i n t o anormalstate;1 --H=5.8T;2--H =5.4;3-H=4.8T;4-H =4.2T;5-H=3.6T;6-H=3.0T;7-H =2.4T;8-H =1.8T; 9 -- H = 1.6 T; 10 -- H = 1.2 T. b -- n -- / -- V - characteristic for a relative normal state

dition d < ~ (T) is not always fulfilled. Therefore, vortices penetrate channels in some places while their motion leads to energy dissipation. The alternate explanation assumes a strict fulfillment of the condition d "~ ~ (7). In his theoretical works, t9'2° Galaiko showed that, in narrow superconducting channels containing current, the transition region can exist at current densities Jca < J < Jc2 where Jcl is the depairing current of Bardeen-Tinkham, that is the depairing thermodynamical current, and Jc2 is the second critical current when the normal state is completely recovered. If one uses the relationship deduced by Galaiko and excludes the sample geometry, one may write Jc2/Jcl = 0.76 [Tc (Tc - T)] y2 Since the superconducting channel is narrow, we may assume that the magnetic field affects only the critical temperature, making it lower. For the Nb-A1 alloy under study, the measured dependence Tc (H) is linear in the field range to 6 T and has a slope of about 0.42 K T -1 . If one assumes that the moment of thermal 'breaking' into the normal state in 1-V characteristics (Fig. 9) corresponds to Jc2 in Galaiko's sense (this, generally speaking, is rather arbitrary assumption) and allows for the fact that the ratio between the measured overall critical current density and the calculated value of Jc2 is the filling (packing) factor, then the sample cross-sectional area occupied by the superconductor may be estimated. Estimation gives the

552

filling factor of superconducting channels on the order of 10 -3 , in agreement with the value found from the electron microscopy data. It is clear that the filling factor can be improved at least by an order of magnitude by optimizing the processes involved in alloy treatment. In this case, the critical current density should be increased by an order of magnitude.

Ways to increase critical current densities

Irradiation effect. It would be only natural to expect that irradiation with high energy particles might be rather favourable for the increase in current paths density. Actually, under definite irradiation conditions, the decomposition of a supersaturated solid solution must be enhanced simultaneously with the stabilization of its dislocation structure, which is produced by the preceeding deformation. Therefore, we studied the electron irradiation effects on decomposition of a gallium supersaturated solid solution in vanadium. The effect of irradiation with 3' particles from 6°Co was also studied but without any interesting results. The alloy V-16 atm% Ga with a 0.8 arm % Ce addition was chosen. Samples had the shape of a ribbon 100 #m thick (the initial state of a deformed solid solution with the bcc phase) and were subjected to the following thermoradiation treatment: irradiation with fast electrons at 40 to 50°C with subsequent isothermal annealing at 700°C. Nonirradiated samples of etalon underwent isothermal annealing

CRYOGENICS. SEPTEMBER 1981

also at 700°C. kradiation was performed in a linear accelerator by 3.5-MeV pulsed electrons at a mean current density of 8 to 10/aA cm -2 . During irradiation the samples were cooled by running water. When deformed samples of a supersaturated solid solution of gallium in vanadium are irradiated with fast electrons, their residual electrical resistance was found to increase. At an electron dose of 1018 el cm -2 this increase was 6.7%. In the investigated range of electron doses, D, the 5R (D) value grows proportionally to D °'6° (Fig. 10). The experimentally observed dependence of the relative electrical resistance on irradiation dose apparently indicates the growth of segregation regions under conditions of nonstationary, radiation-induced diffusion. A smooth run of the electrical resistance temperature dependence at 4.2 to 20 K measured after irradiation of 1018 el cm -2 proves that the electrical resistance at this decomposition stage varies because of segregation, not the precipitation of the V3 Ga A15 phase under irradiation. Isothermal annealing at ?00°C results in the formation of A15 phase precipitates. This may be seen from the measurements of electrical resistance at 4.2 to 20 K after relevant isothermal treatment. Fig. 11 shows that the character of the superconducting transition depends markedly on the preliminary irradiation dose applied to the sample under study. From the comparison of the superconducting transition curves (in a zero magnetic field), it may be seen that electron irradiation before isothermal annealing helps the formation of the structure of superconducting precipitates and gives much lower ATc. This effect becomes more prominent with increasing doses of the preliminary irradiation. Isothermal annealing produced in the process of irradiation results in V3 Ga precipitates with a wider superconducting transition. Kinetics of V3 Ga precipitation from a supersaturated solid solution was investigated by measuring the residual electrical resistance at 77 K. Such investigations have been done most thoroughly on the alloy that had a preliminary irradiation dose of 1018 el cm -2 and on the etalon free from irradiation. As may be seen from Fig. 12, a short (10 to 12 rain) anneal of the irradiated sample increases its electrical resistance from 6.7% (the results of preliminary irradiation) to 13.6%, whereas a 5 min anneal increases the relative electrical resistance of non-irradiated samples to 6.2%. Further

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annealing leads to a smooth decrease in the electrical resistance, this process being faster in the irradiated sample. The alloy with a preliminary irradiation of 10 a el cm-2 displays faster decomposition at early stages (the decomposition time constant almost an order of magnitude smaller than for non-irradiated samples). The reason may lie in the increased number of precipitation centres of V3 Ga phase in the sample that had been irradiated beforehand. Since dislocations serve as precipitation centres, the irradiation may be supposed to stabilize dislocation structure. Conclusions

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CRYOGENICS

. SEPTEMBER

1981

Structurally, multiple connected superconductors prepared by the aging of deformed alloys are no longer uncommon. Provided the relevant technology is developed, such superconductors can compete with traditional multifilamentary superconducting materials. In the development of practical technological processes, three main problems must be solved:

553

one, significant supersaturation o f a solid solution is required to ensure high overall current-carrying capability, and as a rule, this results in poor deformability; two, systems should be chosen where solute atoms segregate on dislocations because this phenomenon ensures elongation, bonding and stoichiometry of the precipitating superconductor phase; and three, the problem of superconductor stabilization b y means of a normal metal with high electrical and thermal conductivity (eg, copper) must be solved either by using a diffusion barrier (eg, a suitably shaped bronze) to avoid diffusion redistribution of the B component or by depositing a stabilizing layer after the aging anneal.

9 10

11

However, the reward, in the sense o f the superconducting and mechanical properties of materials and the advantages o f a simpler production technology, should be great enough to stimulate the solution o f these problems. Structurally, multiple connected superconductors have apparently high potential in the near future.

12

References

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Kaufmann, A.R., Pickett, J.J. Multffilament Nb 3Sn superconducting wire, JAppl Phys 42 1 (1971) 58 Suenaga, M., Sampson, W.B. Superconducting properties of multifilamentary V 3Ga wires, Appl Phys Lett 18 12 (1971) 584-86 Roberge, R., Fihey, J.L. Origin of superconductivity in copperniobium alloys, JAppl Phys 48 3 (1977) 1327-31 ;in: Manufacture of superconducting materials, ed. R.W. Meyerhoff, ASM (1977) 223 Tsuei, C.C. Ductile superconducting copper-base alloys, Science 180 6 (1973) 57-8; Ductile superconducting Cu rich alloys containing A-15 filaments, IEEE Trans Magn MAG-11 (1975) 272-75 Itarbison, J.P., Bevk, J. Superconducting and mechanical properties of in-situ formed multifilamentary Cu-Nb 3 Sn composites, JAppl Phys 48 12 (1977) 5180-87

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Pan, V.M., Beletsky, Yu. I., Flis, V.S., Firstov, S.A., Sarzhan, G.F. Disiokatsionnaya structura, protsessy stareniya i kriticheskiye toki splavov vanadii-haUii, Fiz Met Metalloved 40 2 (1975) 281-88

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Verhoeven, J.D, Finnemore, D.K., Gibson, E.D., Ostenson, J.E., Goodrich, L.F. Superconducting properties of in-situ prepared Nb-Cu-Sn alloys, Appl Phys Lett 33 2 (1978) 101-2 Bormann, R., Schultz, L., Freyhardt, ILC. Superconducting wires, Appl Phys Lett 32 1 (1978) 79-81

B.B., Adams, J., Eager, T.W., Rose, R.M. Superconducting Cu-Nb s Sn composites produced by cold extrusion of fine powders, IEEE TransMagn MAG-15 1 (1979) 689-92 Hemachalam, K., Pickus, M.R. Niobium-tin superconducting wire by the infiltration process, J Less Common Met 46 2 (1976) 297 Pan, V.M., Beletsky, Yu. I. Sverkhprovodniki s nitevidnoy napravlennoy structuroy dispersnoy intermetallidnoy fazy, Sovetsko-Yaponskaya konferentsiya po fizike niskikh temperatur, Tezisy dokladov, Novosibirsk, C-34 (August 14-20, 1969) 30 Pan, V.M., Beletsky, Yu. I., Lasareva, LS., Firstov, S.A. The critical currents in superconductors of the II kind with a disperse intermetalhc phase, Proc LT 12, Kyoto, Japan, Academic Press (1971) 499 Pan, V.M., Beletsky, Yu. I., Sudovtsov, A. I. Issledovaniye svoistv deformiruyemykh sverkhprovodyaschikh splavov vanadii-hallii, Fizika i khimiya obrabotki materialov, N3 (1971) 128-31 Pan, V.M., Beletsky, Yu. I. Sverkhprovodniki 11 roda s dispersnoy intermetallidnoy fazoy;in: Metalloflzika 45 Izdatelstvo 'Naukova dumka', Kiev ( 1973) 81-4 Bardeen, J., Critical fields and currents in superconductors, Low TempPhys ed. C.D. Dewitt et al., New York (1962) 149 Svechnikov, V.N., Pan, V.M., Spektor, A.C. Vliyaniye tseriya na fazoviy sostav i svoistva splavov vanadii-hallii, in: Metalloflzika 27 Izdatelstvo 'Naukova dumka', Kiev (1970) 177

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Pan, V.M., Latysheva, V.I., Khusid, E.N., Minakov, V.H. Turtsevich, E.V., Kriticheskiye toki deformirovannykh stareyuschikh sverkhprovodyaschikh splavov niobii-aluminii, Fiz Met Metalloved 40 3 (1975) 512-17 Pan, V.M., Beletsky, Yu. I., Flis, V.S., Latysheva, V.I., Khusid, E.N., Shpigel, A.S.K.riticheskiye toki deformiruyemykh stareyuschikh sverkhprovodyaschikh splavov vanadii-hallii i niobii-aluminii, in: Sverkhprovodimost, Trudy konferentsii po tekhnicheskomu ispolsovaniyu sverrkhprovodnikov, vol. LV, Sverkhprovodyaschiye materialy, Atomizdat, Moscow (1977) 92-96 Galaiko, V.P. Kritieheskiye toki dlya resistivnykh sostoyanii v sverkhprovodyaschikh kanalakh, Zh Eskp. Teor Fiz 66 1 (1974) 379-85 Galaiko, V.P. O mieroskopicheskoy teorii resistivnykh tokovykh sostoyanii v sverkhprovodyaschikh kanalakh, Zh Eksp Teor Fiz 68 1 (1975) 223-27

CRYOGENICS

. S E P T E M B E R 1981