M.T. Naus, M.C. Jewell, P.J. Lee and D.C. Larbalestier. Applied Superconductivity Center, University of Wisconsin Madison Madison, WI USA

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1 Paper I-09A-10 presented at ICMC-CEC 2001, July 16-20, 2001, Madison, WI USA Published in Advances in Cryogenic Engineering: Proceedings of the International Cryogenic Materials Conference - ICMC, Vol. 48, 48 (B): , 2002, edited by B. Balachandran et al American Institute of Physics This web version for non-commercial educational use only. LACK OF INFLUENCE OF THE Cu Sn MIXING HEAT TREATMENTS ON THE SUPER- CONDUCTING PROPERTIES OF TWO HIGH-Nb, INTERNAL-Sn Nb 3 Sn CONDUCTORS M.T. Naus, M.C. Jewell, P.J. Lee and D.C. Larbalestier Applied Superconductivity Center, University of Wisconsin Madison Madison, WI USA ABSTRACT The heat treatment of internal-sn, Nb 3 Sn superconductors has two primary purposes - (1) to mix the Cu and Sn, and (2) to form the A15 phase. We have studied the effect of the Cu Sn mixing heat treatments on the superconducting properties of two high-nb internal- Sn conductors. The samples underwent continuous ramps of 6, 30 or 60 C/h from room temperature to the A15 reaction temperature, or used isothermal hold treatments to mix the Cu and Sn prior to A15 formation. The strands were then reacted for 180 hours at 650 C to form the A15 phase. Inductive T c measurements showed no significant difference arising from vastly different mixing steps. However, the T c values were ~2 K below that of stoichiometric Nb 3 Sn (18.3 K) due to a high Nb:Sn ratio and Sn leakage from the filament bundle. H*, H c2 and J c (as inferred from magnetization measurements) were also unaffected by the Cu Sn mixing heat treatments used. Unlike lower Nb content, lower J c internal-sn conductors, we conclude that the superconducting properties of these internal- Sn composites are essentially independent of how the Cu and Sn are mixed, greatly simplifying heat treatment optimization. INTRODUCTION The heat treatment (HT) of internal-sn, Nb 3 Sn composites requires the mixing of Sn from the filament bundle core with the interfilamentary Cu in order for each Nb filament to be adequately surrounded by Sn prior to Nb 3 Sn formation. Due to an extremely slow approach to equilibrium, the best mixing heat treatment is complex and varies strongly by composite geometry and/or manufacturer [1]. In principle, the goal of the mixing step is to achieve a homogeneous mixture of Cu and Sn within the Nb diffusion barrier, but in fact it is not possible to fully mix them in a reasonable time (< 200 hours) [2-6]. This inevitable incomplete mixing results in different Nb filaments being in contact with different Cu Sn compositions at the time of A15 formation. From the Cu Sn Nb phase diagram [7], higher Sn A15 is in equilibrium with the higher Sn Cu Sn phases. This implies that the initial A15 to form will have different, 1016

2 perhaps better, superconducting properties because a higher Sn content in the A15 phase should lead to higher T c and H c2. The question then arises whether the presence of different Cu Sn compositions at the start of A15 reaction will have an effect on the A15 properties of fully reacted wires. We have previously reported that there was no discernable effect on the superconducting properties in wires that underwent various sub-600 C hold temperatures for the Cu Sn mixing HT when followed by a reaction HT of 180 hours at 650 C (180 h/650 C) [6]. However, the full range of possible mixing HT was not explored, as the lowest and highest Sn mixing temperatures were 362 and 533 C, respectively. In response to concerns about the generality of this earlier result, we have investigated the more standard (and lengthy) manufacturer s HT and compare results to greatly abbreviated Cu Sn mixing heat treatments. In this paper, we test a full range of possible Cu Sn mixing heat treatments. We investigate the effect on the transition temperature (T c ), the irreversibility field (H*), and the upper critical field (H c2 ) of continuous 6, 30 and 60 C/h ramp rates from room temperature to a reaction temperature of 650 C. We compare their properties to those of samples inserted directly into a furnace at 650 C and to those that underwent a more standard, long Cu Sn mixing HT step. We also compare the critical current density (J c ) (as inferred from magnetization measurements) for a sub-set of these heat treatments. EXPERIMENTAL PROCEDURE The wires in this study (FIGURE 1), designated CRe1912 and ORe102, were manufactured by the modified jelly roll[8], internal-sn process [9] by Teledyne Wah Chang (now Wah Chang) and Oxford Instruments Superconducting Technology, respectively. Their J c (4.2 K, 12 T) values using a resistivity criterion of Ω m are 2200 A/mm 2 [10] and 2400 A/mm 2 [11] for CRe1912 and ORe102, respectively. Both wires have 54 bundles, each bundle consisting of an alloyed Sn core surrounded by alternating layers of Cu sheet and Nb-1wt.%Ti mesh. CRe1912 has a Sn-2wt.%Mg core, while ORE102 has a Sn-1wt.%Cu core. The wire diameters are 0.7 and 0.8 mm, respectively. Around each bundle is a Nb diffusion barrier. TABLE 1 shows the Cu:Nb:Sn atomic ratios derived from SEM image analysis on polished cross-sections like those in FIGURE 1. The reacted microstructures showed considerable conversion of the Nb diffusion barrier around each bundle. Thus, TABLE 1 shows the Cu:Nb:Sn ratios both with and without the barrier included in the calculation. A potentially significant design difference between CRe1912 and ORe102 is that the atomic Nb:Sn ratios before reaction were 3.2 and 2.6, respectively, making CRe1912 Sn-poor and ORe102 Sn-rich. When including the barrier, these ratios increased to 4.4:1 and 3.6:1 for ORe102 and CRe1912, respectively, making both conductors strongly sub-stoichiometric. The ends of ~80 mm long samples were electroplated with Cu to prevent Sn leakage. Samples were cleaned and sealed in quartz tubes under ~30 mtorr of Ar, then placed in a furnace at room temperature, heated to 650 C at 6 C/h, 30 C/h or 60 C/h and reacted for 180 hours at 650 C (180 h/650 C). Another sample set was placed directly into a pre- TABLE 1. Cu:Nb:Sn Ratios Prior to Reaction Wire Designation Including Entire Barrier Not Including Barrier Cu Nb Sn Comments Cu Nb Sn Comments CRe Sn deficient Sn deficient ORe Sn deficient Sn rich

3 (a) (b) Nb-1wt%Ti Filaments Cu-Mg-Sn Nb-1wt%Ti Filaments Voids Sn-2wt%Mg Core Nb diffusion barrier Sn-1wt%Cu Core Nb diffusion barrier FIGURE 1. Backscatter scanning electron microscope (SEM) images of (a) CRe1912 and (b) ORe102 prior to reaction. heated furnace at 650 C for 180 hours to form the Nb Sn A15 phase. It is assumed that these samples reached 650 C within a few minutes. Comparison of these simple heat treatments was made to a complex but standard manufacturer s heat treatment of 120 h/185 C (solid state mixing below the melting point of Sn) + 72 h/340 C (solid state mixing below the formation temperature of δ phase Cu Sn) h/650 C, with a ramp rate of 60 C/h between hold temperatures. Examination by light microscopy revealed that a HT of 104h/340 C generated the same microstructure as 120 h/185 C + 72h/340 C. Therefore, a heat treatment of schedule of 104h/340 C h/650 C was also performed. After reaction, approximately 10 mm of each sample was removed to avoid any effects of Sn diffusion into the Cu-plated ends. Inductive T c measurements were conducted with a Quantum Design SQUID magnetometer. The ~3 mm long samples were zero-field cooled to 6 K, 5 mt was applied and their moment measured upon warming with the wire axis parallel to the applied field. H* and H c2 were measured at 12 K and 4.2 K using an Oxford Instruments 14 T vibrating sample magnetometer (VSM) with the ~5 mm long samples oriented normal to the applied field. From the Bean model [12], J c is proportional to the hysteretic loop width ( m) and, as we did not have the A15 volume measurement to calculate J c, Kramer plots [13] were generated by plotting m 1/2 B 1/4 vs. B. H* was then determined by extrapolation of the Kramer plot to zero (H* Kramer ). H c2 was determined as the field at which the sample moment became less than the paramagnetic background moment. 1018

4 Normalized Moment (a.u.) CRe1912 ORe102 6 C/h 30 C/h 60 C/h 6 C/h 30 C/h 60 C/h (a) Temperature (K) Normalized Moment (a.u.) CRe1912 ORe102 Full HT CRe1912 Full HT ORe102 (b) Temperature (K) FIGURE 2. T c of ORe102 (open symbols) and CRe1912 (closed symbols) for various heat treatments. In (a), the samples were ramped from room temperature to 650 C at the ramp rates listed or inserted directly into a pre-heated furnace, and then held for 180 hours at 650 C. No trend was seen in the T c as a function of ramp rate. A comparison of a standard manufacturer s HT ( Full HT ) to one in which the sample was inserted directly at 650 C ( ) is shown in (b). The difference in T c was ~ 0.4 K. Both composites have broad transitions with mid-points near 16 K, well below the T c of stoichiometric Nb 3 Sn. m was normalized by sample mass for ORe102 to get a relative comparison of J c values after 3 different heat treatments. The heat treatments examined in this way were (1) direct insertion for 180 h/650 C, (2) 104 h/340 C h/650 C, and (3) 120 h/185 C + 72 h/340 C h/650 C. RESULTS The T c curves are shown in FIGURE 2. As a whole they are broad and significantly lower than expected for high-nb conductors. Their mid-point T c transitions occur at ~15.8 K for CRe1912 and at ~16.3 K for ORe102, and their 10% to 90% transitions were ~1 K wide. There was no discernible difference in the T c values for different ramp rates (Figure 2a), although ORe102 always had a slightly higher T c than CRe1912. Figure 2b shows the T c traces after a standard manufacturer s HT of 120 h/185 C + 72h/340 C h/650 C ( Full HT ) and after putting the samples directly into a 650 C furnace for 180 hours ( ). The difference between the two heat treatments is less than 0.4 K. The T c traces of the 104 h/340 C h/650 C HT lay directly over the Full HT curves and, for clarity, they were not plotted. TABLE 2 shows H c2 (12 K) ~13.3 T, H* Kramer (12 K) ~10.0 T and H* Kramer (4.2 K) ~24.2 T after all heat treatments for CRe1912. In a few instances H c2 could not be accurately defined due to curvature in the magnetization curve and these are represented by an asterisk (*) in the table. The measurement errors for H* Kramer (12 K) and H c2 (12 K) are estimated to be 0.1 T and 0.2 T, respectively. The errors for H* Kramer (4.2 K) are unknown because we must extrapolate from 14 T to ~24 T, but we expect them to be small because there is good consistency in the results. TABLE 3 shows that H* and H c2 for ORe102 are also not significantly different for the different heat treatments. H c2 (12 K), H* Kramer (12 K) and H* Kramer (4.2 K) were ~13.3 T, ~10.0 T and ~24.4 T, respectively. Also shown are three hysteretic loop widths normalized to sample mass ( m/mass), which includes the longest and shortest Cu Sn 1019

5 TABLE 2. H* Kramer and H c2 of CRe1912 H* Kramer (T) H c2 (T) Heat Treatment 4.2 K 12 K 12 K 120 h/185 C + 72h/340 C h/650 C * 104h/340 C h/650 C C/h, RT to 180 h/650 C C/h, RT to 180 h/650 C * 60 C/h, RT to 180 h/650 C insertion, RT to 180 h/650 C * * = Unable to determine. mixing heat treatments. The values at 4.2 K & 12 T and at 12 K & 5 T differ by less than 1%. It is interesting to note that H* Kramer (4.2 K) values for ORe102 are consistently slightly below those of CRe1912, even though ORe102 has a slightly higher T c. However, the H* Kramer (12 K) values are essentially the same for the two wires. DISCUSSION TABLE 2 and 3 clearly show that the Cu Sn mixing HT has no influence on the superconducting properties in high-nb wires. This is true whether the mixing HT lasts 200 hours or just a couple of minutes (i.e. direct insertion). This is unlike lower Nb composites (e.g. ITER type), where the aggressiveness of the Cu Sn mixing HT influences the Cu Sn phase boundary movement and thereby the amount of filament coupling [14]. However, this is not an issue for high-nb wires because the filaments are initially so close together that it is not possible to keep them from coupling during heat treatment. The T c of these high-nb conductors is significantly lower than expected. The midpoint transitions are ~16.3 and ~15.8 K for ORe102 and CRe1912, respectively, ~2 K lower than that of stoichiometric Nb 3 Sn (18.3 K). A possible cause of the suppressed T c values is compressive strain induced by the higher thermal contraction of Cu relative to that of Nb 3 Sn. As a test, the Cu was etched off a fully reacted CRe1912 sample discussed in [6]. The inductive T c of a single, free bundle was measured to be only 0.5 K higher than that in the unetched, strained state. Therefore, we conclude that the primary cause of the suppressed T c values of the wires in this study cannot be due to strain and is most likely an A15 composition effect. A consequence of the high-nb design is that the Nb barrier takes part in the reaction, decreasing the average A15 Sn concentration and passing Sn into the stabilizing Cu, away from the filaments. If the barrier reacts significantly, both wires will be Sn deficient (see TABLE 3. H* Kramer, H c2 and m/mass of ORe102 H* Kramer (T) H c2 (T) m/mass (emu/mg) Heat Treatment 4.2 K 12 K 12 K 4.2 K, 12 T 12 K, 5 T 120 h/185 C + 72h/340 C h/650 C h/340 C h/650 C C/h, RT to 180 h/650 C * C/h, RT to 180 h/650 C C/h, RT to 180 h/650 C insertion, RT to 180 h/650 C * = Unable to determine. -- = Not measured. 1020

6 Cu Stabilizer Nb Barrier A15 grains Compromised Barrier FIGURE 3. Backscatter SEM image of ORe102 after direct insertion at 650 C for 180 hours showing the region between 2 bundles where the A15 phase has grown completely through the barrier. As a result of Sn leakage into the Cu stabilizer, A15 grains can be seen growing on the external side of the Nb barrier. TABLE 1), with CRe1912 more Sn deficient than ORe102. This is consistent with the fact that ORe102 has a higher T c than CRe1912. Clear evidence for significant barrier reaction is seen in the backscatter SEM image shown in FIGURE 3 of an ORe102 barrier region after direct insertion into a furnace for 180 h/650 C. The image also shows a region where the barrier has reacted fully through. There also appear to be A15 grains on the external side of the Nb barrier, indicating Sn leakage into the stabilizing Cu. As further proof of contamination of the pure stabilizing Cu, the resistivity ratio (ρ 300 K /ρ 77 K) was only 3.1. Suppressed RRR values were also seen in the internal-sn wires examined by Barzi et al. [15]. As a result, the Sn deficiency is now exacerbated by Sn loss into the stabilizing Cu. Moreover, some residual Sn will be left in the interfilamentary Cu after reaction. Therefore, both wires may be even more Sn deficient than their pre-reaction Cu:Nb:Sn ratios would indicate. Although the above discussion provides a self-consistent explanation of why T c is reduced for both conductors, and why T c is higher in ORe102, the H* Kramer and H c2 data is less clearly explained. H* Kramer (4.2 K) values for the higher T c ORe102 (23.2 to 24.0 T) are consistently lower than those of CRe1912 (24.1 to 24.3 T). Since H* Kramer may be controlled by grain size (the Mg in CRe1912 is an A15 grain refiner [16]), it may be that the small differences of TABLE 2 and 3 are explained by these rather than the overall Sn content. Also, the H c2 detected is the highest in the wire, but it is not necessarily that of the majority of the A15 material. Even though there is no effect on the Nb 3 Sn superconducting properties, an additional consideration that often forces serious concern about the Cu Sn mixing HT is that of Snburst, in which liquid Sn ruptures the wire. This can happen when the HT temperature exceeds the melting point of Sn (232 C) before the Sn and Cu have had a chance to form an adequate amount of higher melting point Cu Sn phases. Although we have not seen Snburst in this study, the probability of it occurring increases with sample length due to the increased probability of encountering weak diffusion barrier regions. The Sn-burst probability might also be increased if cabling of the wire sufficiently reduces the integrity of the barrier. 1021

7 CONCLUSION We have looked at the effect of vastly different Cu Sn mixing heat treatments on the superconducting properties of fully reacted high Nb, internal-sn Nb 3 Sn wire. We can find no significant influence on T c, H* Kramer, H c2 or J c due to the Cu Sn mixing heat treatments. Since all of the filaments couple in these high-nb wires, attention only needs to be paid to the A15 reaction step. However, the integrity of the Nb barrier in these designs is at significant risk. Barrier breach and stabilization Cu poisoning is possible even likely leading to reduced T c, H* Kramer and H c2 values in the A15 layer. Despite these concerns, high J c values (defined as the critical current normalized to the bundle cross-section) have been achieved. Clearly, there presently exists a tradeoff between the area cross-section of the A15 layer and the A15 superconducting properties. If high-nb, internal-sn wires are to reach their full potential, the issues of the proper Nb:Sn ratio and Sn leakage into the stabilizing Cu must be addressed. ACKNOWLEDGEMENTS This work was supported by the U.S. Department of Energy Office of Fusion Energy Sciences (DE-FG02-86ER52131) and Division of High Energy Physics (DE-FG02-91ER40643), and benefited from NSF-MRSEC (DMR ) supported facilities. The authors would like to thank Oxford Instruments - Superconducting Technology and Wah Chang for donating the wire in this study. REFERENCES 1. Takayasu, M., Childs, R.A., Randall, R.N., Jayakumar, R.J. and Minervini, J.V., IEEE Trans. Appl. Supercond., 9 (2), pp (1999). 2. Dietderich, D.R., Glazer, J., Lea, C., Hassenzahl, W.V. and Morris, J.W. Jr., IEEE Trans. Mag., 21 (2), pp (1985). 3. Taillard, R. and Verwaerde, C., Proc. of the 4 th European Conference on Advanced Materials and Processes (Euromat 95), Associazione Italiana di Metallurgia, Milan, Italy, 1995, pp Glowacki, B.A., IEEE Trans. Appl. Supercond., 7 (2), pp (1997). 5. Naus, M.T., Lee, P.J. and Larbalestier, D.C., IEEE Trans. Appl. Supercond., 10 (1), pp (2000). 6. Naus, M.T., Lee, P.J. and Larbalestier, D.C., IEEE Trans. Appl. Supercond., 11 (1), pp (2001). 7. Villars, P., Prince, A. and Okamoto, H., Handbook of Ternary Alloy Phase Diagrams, ASM International, 1995, p McDonald, W.K., Curtis, C.W., Scanlan, R.M., Larbalestier, D.C., Marken, K. and Smathers, D.B., IEEE Trans. Mag., 19 (3), pp (1983). 9. Hashimoto, Y., Yoshizaki, K. and Tanka, M., Processing and properties of superconducting Nb 3 Sn filamentary wires, in Proc. of the 5 th International Cryogenic Engineering Conference, edited by K. Mendelssohn, IPC Business Press Ltd., London, 1974, pp Zhang, Y., McKinnell, J.C., Hentges, R.W. and Hong, S., IEEE Trans. Appl. Supercond., 9 (1), pp (1999). 11. Personal communication, Michael B. Field, Oxford Instruments Superconducting Technology. 12. Bean, C.P., Rev. Mod. Phys., 36, pp (1964). 13. Kramer. E.J., J. Appl. Phys., 44 (3), pp (1973). 14. Suenaga, M. and Sabatini, R.L., Effects of Temperature Ramp Rate during Heat Treatment on Hysteresis Loss and Critical Current Density, in Proc. of the 9 th US-Japan Workshop on High Field Materials, edited by K. Osamura et al., 1995, pp Barzi, E., Limon, P.J., Yamada, R. and Zlobin, A.V., IEEE Trans. Appl. Supercond., 11 (1), pp (2001). 16. Togano, K., Asano, T. and Tachikawa, K., J. Less Comm. Met., 68, pp (1979). 1022