Cryogenics 103 (2019) 102974
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Comparison of phase transformation and superconducting properties of RHQT Nb3Al wires fabricated at Tc peak and valley condition
T
⁎
Lian Xiaa, Xiaguang Suna, Pingyuan Lia, Zhou Yua, , Yongliang Chena, Yong Zhanga, Xifeng Panb, ⁎ Guo Yanc, Yong Fengc, Yong Zhaoa,b, a
Key Laboratory of Advanced Technology of Materials (Ministry of Education), Superconductivity and New Energy R&D Center (SNERDC), Southwest Jiaotong University, Chengdu, Sichuan 610031,China b College of Physics and Energy, Fujian Normal University, Fuzhou, Fujian 350117, China c National Engineering Laboratory for Superconducting Materials (NELSM), Western Superconducting Technologies (WST) Co., Ltd, Xi’an 710018, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Nb3Al wires Rapid heating Phase transformation Annealing temperature Grain boundary pinning
This work compared the phase transformation feature and superconducting properties of Nb3Al wires rapidly heated at 272 A and 315 A, closing to the Tc peak and valley current condition, respectively. The fabricated rapidly heating, quenching and transformation (RHQT) Nb3Al wires are mainly consisted of A15 phase with trace Nb and Nb2Al impurity phase. With the increase of annealing temperature from 700 °C to 1100 °C, the Tc value and Jc performance of Nb3Al wires improve first and then gradually decrease. M-T and Jc-H curves reveal that the best Tc and Jc performance of 272 A samples was obtained at the annealing temperature of 800 °C, compared to that of 900 °C for the 315 A samples. In both 272 A and 315 A samples, lower annealing temperature cannot promote the phase transform process from bcc to A15 phase. Excessive heating beyond 1000 °C caused the already transformed Nb3Al phase to decompose into Nb2Al and Al-poor Nb3Al phase, resulting in degraded Tc and Jc performance of the Nb3Al wires. The main pinning mechanism of Nb3Al wires was grain boundary pinning, as deduced from the fitting of flux pinning force versus applied field curves.
1. Introduction Compared to Nb3Sn, Nb3Al have higher superconducting transition temperature (Tc), larger upper critical field (Hc2) [1,2], and better tolerance to applied strain [3,4]. The magnets wound using Nb3Al superconducting wires can transport stronger current under higher background field. Nb3Al wires was suitable application in the fields of particle accelerator magnet or fusion devices [1]. Rapid heating, quenching and transformation (RHQT) technique has been developed for decades to fabricate long length Nb3Al wires with excellent superconducting performance [5,6]. Risk of wire breakage during reel to reel RHQ process because of the non-uniformity of wire, including diameter, resistance and strength requires precisely control of apparatus parameters, including fluctuation of heating distance, vibration and tension force of the wires [7]. Since RHQT Nb3Al wire is fabricated by a phase transformation process, the critical current density (Jc) properties of Nb3Al wires are principally controlled by the rapid heating and quenching or transformation conditions. During the RHQ process, a direct current passed through the precursor Nb-Al composite wire and rapidly Ohmic-heated the wires to
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beyond 1900 °C in less than 1 s. Then the wire was quickly quenched in a liquid Ga bath. The resulted microstructure and superconducting properties of Nb3Al wire strongly depends on the applied current of RHQ (IRHQ) which directly determines the Tmax of the wire. With the increase of IRHQ, the finally transformed Nb3Al wires can be divided into four regions [7]: (1) A15 and Nb2Al phases, (2) extended solid solution, (3) partial melting, and (4) complete melting. For the Nb3Al wire quenched from body centered cubic (bcc) extended solid solution region where the Tmax reaches between 1940 and 2060 °C, the scatter of resultant Tc and Jc of the transformed Nb3Al wire with the IRHQ is almost negligible. When the wire was rapidly heated at larger IRHQ entering the partial melting region, Jc of the fabricated Nb3Al wires dynamically fall down because of formation of Al-poor solid solution and Al-rich liquid phases, leading to off-stoichiometry composition of the fabricated Nb3Al [7]. Concern the effect of post heat treatment, Banno et al. [8] had identified that the transformation temperature from bcc to A15 phase is between 730 °C and 800 °C. They also developed unique process of transformation-heat-based up-quenching (TRUQ) to improve the stoichiometry of the Nb3Al and enhanced the Jc of the Nb3Al superconductors beyond 25 T [9]. For the Nb3Al wires annealed at 800 °C for
Corresponding authors. E-mail addresses:
[email protected] (Z. Yu),
[email protected],
[email protected] (Y. Zhao).
https://doi.org/10.1016/j.cryogenics.2019.102974 Received 15 February 2019; Received in revised form 6 August 2019; Accepted 30 August 2019 Available online 31 August 2019 0011-2275/ © 2019 Elsevier Ltd. All rights reserved.
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the magnetization hysteresis loops (M-H) were measured by a physical property measurement system (MPMS, Quantum Design). The critical current density Jc = 15ΔM/(V × R) was calculated from the M-H curves using the Bean model, where the ΔM is the difference between the positive and negative magnetic moments corresponding to the hysteresis loop of M-H in the process of ascending and descending magnetic fields, V is the volume of the test sample and R is the radius of the cylindrical wire.
10 h, a peak in the Tc-IRHQ curve appears before the Tc and Jc reaches the plateau region which corresponds to the optimal stable RHQ condition for long length wire fabrication. When the wire is quenched from the partial melting region, a valley of Tc and Bc2 was observed in the RHQT Nb3Al wires [7]. In a sense, the peak and valley of Tc can be regarded as important marks indicating the boundary of the optimism RHQ condition. Thus, it is interesting to systemically study the properties of RHQT Nb3Al prepared at the Tc peak and valley region, to clarify the mechanism behind their unique superconducting properties. In this paper, we compared the phase transformation feature and superconducting properties of Nb3Al wires rapidly heated around the Tc “peak” and “valley” condition. The results show that Tc and Jc of Nb3Al wires reaches the maximum value at different transformation temperature related to their initial quenched state. When the IRHQ is around the Tc peak current condition, the temperature to obtain the best Tc and Jc of the Nb3Al wires is lower, at 800 °C, but when the IRHQ is around the Tc valley, the optimism temperature becomes higher at 900 °C.
3. Results and discussion Fig. 2 shows the XRD patterns of the Nb3Al wires. Wires annealed at 700 °C are consisted of Nb phase. For the samples annealed beyond 800 °C, Fig. 2(a) and (b) reveal that the wires are mainly composed of Nb3Al A15 phase. Besides Nb3Al phase, trace Nb2Al phase was detected in the annealed wires, which might attributed to the insufficient reaction of Nb and Al element because of inhomogeneous distribution of Nb and Al layers during wire cold working process because of their significant hardness difference. Elevation of annealing temperature leads to increase of diffraction peak intensity of Nb3Al phase, indicating crystallinity improvement of Nb3Al phase. Nb was found in the wires since Nb presented at the center of precursor wires. According to previous report [16,17], 272 A close to the Tc peak current condition, and the 315 A wire corresponds to that of obtaining Tc valley. Fig. 3 illustrates the temperature dependence of DC magnetization for the annealed samples. The onset transition temperature, Tc-onset has been illustrated by inset of Fig. 3(a) and (b). The amplitudes of the diamagnetic signals were comparable for the 315 A and 272 A samples, indicating almost same volume ratio of A15 Nb3Al phase in the samples. Fig. 4(a) shows the Tc-onset value of the fabricated Nb3Al wires. The Tconset of the wires increases first and then decreases with the increase of annealing temperatures from 700 °C to 1100 °C. For the 272 A RHQ samples, the Tc-onset value increases from 14.5 K at 700 °C to 16.5 K at 800 °C and then monotonously decreases to 13.8 K at 1100 °C. For the 315 A rapidly heated samples, wire annealed at 700 °C shows Tc-onset value of only 9.3 K, which is corresponding to the superconducting transition of Nb. 315 A sample annealed at 900 °C exhibits a significant increase of Tc-onset value to 16.5 K, indicating complete phase transition of Nb(Al)ss → A15 phase. Tc-onset value decrease to 15.6 K and 15 K when the 315 A sample annealed at 1000 °C and 1100 °C, which are almost similar to that of 272 A sample annealed at 900 °C and 1000 °C. Jorda et al. [17] found that Tc of Nb3Al increases with the improvement of Al content in Nb3Al. Hence, the possible reason for the decrease tendency of Tc is the deviation from the stoichiometry composition of the Nb3Al. Fig. 4(b) shows the transition width (ΔTc) of 272 A and 315 A samples annealed at different temperatures. The ΔTc values are wide, which might attribute to the presence of impurity Nb, Nb2Al and other weak superconducting phases in the formed Nb3Al superconductors. Fig. 5 illustrates the Jc-B curves of the 272 A and 315 A samples annealed at different temperatures. Fig. 5(a) and (b) show that the Jc values of 700 °C annealed 272 A sample decay very fast at higher applied fields, which indicates that thermal energy at 700 °C is insufficient to promote the nucleation and growth of Nb3Al phase [8]. 272 A sample annealed at 800 °C shows significant improvement of Jc at high field and gets the best Jc value of 3.22 × 104 A/cm2 @ 5 K, 6 T. For the 272 A samples annealed at temperatures beyond 1000 °C, the degradation of Jc value at high field become serious again, attributing to the decomposition of the transformed Nb3Al phase to Nb2Al phase and Al-poor Nb3Al phase with degraded superconducting performance. The Jc-B curves of the annealed 315 A samples were shown in Fig. 5(c) and (d), showing similar variation tendency of Jc value as these of 272 A samples. 315 A sample annealed at 900 °C shows best Jc value of 4.1 × 103 A/cm2 @5K, 6 T. Fig. 5(e) illustrates the Jc-B curves at 5 K, 8 K and 10 K of 272 A and 315 A sample annealed at their optimism temperature of 800 °C and 900 °C, respectively. The 272 A wire show
2. Experimental Single core Nb3Al precursor wires was made by Jelly Roll technology [10-12]. Nb foil and Al foil with thickness of 100 μm and 30 μm were wound on a Nb rod, then inserted into a Cu can. This composite was extruded by room temperature hydrostatic extrusion. Then, the wire was cold drawn to final diameter of 1.0 mm. The Cu sheath was etched off before performing the rapid heating and quenching (RHQ) process. The length of Nb/Al precursor wire is several hundred meters. Fig. 1 shows the schematic of our home made RHQ equipment. No tension force was applied to the wires and the parameter of the equipment was carefully adjusted to minimize the wire vibration during reel-to-reel RHQ process. RHQ was carried out under pressure of 10-3 Pa. The precursor wire, moving with a speed of 0.3 m/sec, was continuously heated up to about 2000 °C [13,14] by Ohmic-heating during 0.4 sec, with a DC current transported between a Cu capstan and a molten Ga bath. The wire was subsequently quenched into the molten Ga bath at about 50 °C [15]. The Nb/Al precursor wire was converted into the Nb-Al bcc supersaturated solid solution wire by RHQ at different currents of 272 A and 315 A, roughly corresponding to the temperature of 1900 °C and 2300 °C. The fabricated Nb(Al)ss wires were cut and sealed in quartz tubes. Additional annealing was performed to transform the bcc phase to Nb3Al A15 phase. The annealing temperature was varied from 700 °C to 1100 °C with step of 100 °C. The annealing time was kept constant of 10 h. After annealing, the samples were naturally cooled down to room temperature. Crystalline phase of the annealed samples were analyzed by X-ray diffraction (XRD, Philips X’Pert MRD) using Cu Kα radiation (λ = 0.154 nm). Surface morphology and microstructure of Nb3Al superconductors were investigated by Field-emission scanning electron microscopy (FESEM, Quanta FEG 250). Superconducting transition temperature (Tc) was determined by the M-T curves measured under a magnetic field of 20 Oe with the zero-field cooling (ZFC) pattern and
Fig. 1. Diagram of continuously rapid-quenching system. 2
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Fig. 5. The Jc-B curves of different temperature annealed 272 A samples measured at (a) 5 K and (b) 10 K, Jc-B curves of different temperature annealed 315 A samples measured at (c) 5 K and (d) 10 K, (e) Jc-B curvesof 800 °C annealed 272 A and 900 °C annealed 315 A sample measured at 5 k, 8 K and 10 K; (f) The Gibbs free energy of Nb(Al)ss phase and A15 phase vs Al concentration [11].
272 A and 315 A samples was analyzed as follows. The Gibbs free energy change, ΔG, when the spherical shape particle of A15 phase formed from the bcc matrix phase, is [9],
much better Jc performance than that of the 315 A wire, where Jc reaches 2.3 × 103 A/cm2 and 2.9 × 102 A/cm2 at 10 K, 5 T, respectively. Rapidly heated at 272 A means the sample was treated close to the current condition of getting Tc peak, which is about 1900 °C, and samples rapidly heated at 315 A reaches partial melting. Fukuzaki et al. [18] suggested that excess heating leads to (1) the quenched Nb(Al)ss consisted of off stoichiometry Al-poor bcc and Al-rich bcc and (2) the inevitable reaction between Nb rod and the liquid Nb-Al phase. In this case, higher annealing temperature is required to start Nb3Al A15 nucleation from the Nb(Al)ss matrix [8]. The nucleation process and phase transformation mechanism of the
ΔG = −4 π r 3Δg/3 + 4πr 2σ
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Fig. 7. Typical SEM images of 272 A rapidly heated samples annealed at (a) 800 °C/10 h, (b) 900 °C/10 h, (c) 1000 °C/10 h, (d) 1100 °C/10 h.
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temperature, Tc of RHQT Nb3Al wires increases first and then gradually decreases, indicating change of stoichiometry of Nb3Al A15 phase. For the samples rapidly heated at 272 A, the best Tc and Jc value was obtained at annealing temperature of 800 °C, however, higher annealing temperature of 900 °C is required for samples rapidly heated at 315 A. The initial state of the Nb(Al)ss phase was different for 272 A and 315 A RHQ samples, resulting in difference of annealing temperature to complete the phase transformation process. On one hand, too low annealing temperature cannot promote the phase transform from bcc supersaturated solid solution to A15 superconducting phase. On the other hand, samples annealed at excess temperatures cause the already transformed Nb3Al phase to decompose into Nb2Al phase and Al-poor Nb3Al phase which presents degraded Tc and Jc performance. Fitting of the flux pinning force and field curve reveal that flux pinning mechanism is mainly the grain boundary pinning in the Nb3Al superconducting wires.
of newly formed particle exceed the critical size, the new phase particle will continuous growth and finally form A15 grain. For the 272 A sample, the composition of Al in the formed Nb(Al)ss close to stoichiometry of Nb3Al. Gibbs free energy change of the phase transformation from Nb(Al)ss to A15 phase is shown as arrow I in Fig. 5(f). Rapidly heated at 315 A resulted in formation of Nb(Al)ss with Al poor composition which increase the Gibbs free energy of the quenched Nb(Al)ss, as shown arrow II in Fig. 5(f). Thus the chemical potential is decreased and the free energy barrier is increased during the phase transformation process from Nb(Al)ss to A15. Hence increase of the transformation temperature was expected as a result of the increase of free energy barrier which hinder the nucleation process of the A15 phase. To clarify the pinning mechanism of the Nb3Al superconducting wires, the field dependence of pinning force (Fp) at 10 K was drawn in Fig. 6(a) and (b), where the Fp was calculated by using the formula: Fp = μoH × Jc. With the increase of annealing temperature, the value of maximum flux pinning force (Fp,max) of 272 A and 315 A samples show similar variation tendency of improvement first and then gradually degradation. Fig. 6(a) and (b) show that the Fp,max is about 4.2 × 108N/ m3 and 1.8 × 108N/m3 at applied magnetic field of 1.85 T and 1.45 T for the 800 °C annealed 272 A and 900 °C annealed 315 A sample, respectively. Fitting the experimental curves with Kramer’s scaling law [19] was shown in Fig. 6(c) and (d). The trend of normalized pinning force density (fp = Fp/Fp,max) versus normalized magnetic field (B/Birr) was basically the same under different temperatures. The fitting parameters p and q are about 0.7 and 1.7, respectively, which were close to the parameters of surface pinning model [18], indicating that the primary flux pinning source was derived from grain boundaries. Fig. 7 shows the typical SEM images of the 272 A samples annealed at different temperatures. When annealing temperature is 800 °C and 900 °C, the SEM images exhibit uniform microstructure of the A15 phase. The morphology of 800 °C sample shows unambiguous region with size of about 10 μm where the boundary shows darker contrast in the SEM image. This unambiguous region is consisted of much smaller grains with scattered distribution of grain size. At higher temperature of 900 °C, the dimension of the region shrink and the consisted smaller grains cannot be observed in the SEM images. Excessive annealing (temperature > °C) causes a two-phase microstructure with different contrast. Morphology of Nb phase with light color embedded in the matrix of A15 phase was observed in Fig. 7(c) of 1000 °C annealed sample. This indicates that in some part of the precursor wire, the reaction of the Nb/Al composite proceeded inadequately with the diffusion length of Al layer larger than 1 μm. Therefore, the resultant A15 phases in the Nb3Al wires seems to have a wide composition distribution, resulting in the A15 phases of various Tc. For the 1100 °C annealed samples, voids was formed in the samples. Since bcc supersaturated solid solution Nb(Al)ss phase was formed by rapid heating and quenching from temperatures above 1900 °C, the phase transition from metastable Nb(Al)ss → A15 phase occurs by a massive transformation at low annealing temperature, which requires only a change in crystal structure and not a change in local composition [11]. However, at the composition of Nb–25 at.%Al, a two-phase structure consisted of Nb2Al and Al-poor Nb3Al phase exhibits smaller free energy. When the Nb (Al)ss wires annealed beyond 1000 °C, the heating energy is high enough to promote the long length diffusion of Al atoms, thus the transition Nb(Al)ss → A15 + Nb2Al starts and results in the microstructure of formed Nb2Al particles dispersed in an off-stoichiometric A15 phase matrix.
Declaration of Competing Interest We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work, there is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript submitted to《Cryogenics》, which is entitled “Comparison of phase transformation and superconducting properties of RHQT Nb3Al wires fabricated at Tc peak and valley condition”. Acknowledgements The authors are grateful to the financial support of the National Key R&D Program of China (No. 2017YFE0301401), the National Natural Science Foundation of China (Grant Nos. 61404109, 51377138), and the Program of International S&T Cooperation of China (No. 2013DFA51050). References [1] Takeuchi T, Kikuchi A, Banno N, Kitaguchi H, Iijima Y, Tagawa K, et al. Status and perspective of the Nb3Al development. Cryogenics 2000;48:371–80. [2] Parrell JA, Zhang YZ, Field MB, Cisek P, Hong S. High field Nb3Sn conductor development at Oxford Superconducting Technology. IEEE Trans Appl Supercond 2003;13:3470–3. [3] Specking W, Kiesel H, Nakajima H, Ando T, Tsuji H, Yamada Y, et al. First results of strain effects on Ic of Nb3Al cable in conduit fusion superconductors. IEEE Trans Appl Supercond 1993;3:1342–5. [4] Takeuchi T, Iijima Y, Inoue K, Wada H, ten Haken B, ten Kate HHJ, et al. Strain effects in Nb3Al multifilamentary conductors prepared by phase transformation from bcc supersaturated-solid solution. Appl Phys Lett 1997;71:122–4. [5] Kikuchi A, Iijima Y, Inoue K. Microstructures of Rapidly-Heated/Quenched and Transformed Nb3Al Multifilamentary Superconducting Wires. IEEE Trans Appl Supercond 2001;11:3615–8. [6] Takeuchi T. Nb3Al conductors-rapid-heating, quenching and transformation process. IEEE Trans Appl Supercond 2000;10:1016–21. [7] Fukuzaki T, Takeuchi T, Banno N, Tagawa K, Tatsumi N, Ogiwara H, et al. Effect of rapid heating, quenching and transformation conditions on the superconducting properties and microstructure of Jelly-Roll processed Nb3Al superconductors. Supercond Sci Technol 2002;15:1404–9. [8] Banno N, Takeuchi T, Fukuzaki T, Kitaguchi H, Tagawa K, Iijima Y, et al. Relationship between BCC-Deformation, transformation temperature and microstructure in Nb3Al wires. IEEE Trans Appl Supercond 2007;17:2688–92. [9] Banno N, Takeuchi T, Fukuzaki T, Wada H. Optimization of the TRUQ (Transformation-heat-based up-quenching) method for Nb3Al superconductors. Supercond Sci Technol 2002;15:519–25. [10] Pan XF, Feng Y, Yan G, Cui LJ, Chen C, Zhang Y, et al. Manufacture, electromagnetic properties and microstructure of an 18-filament jelly-roll Nb3Al superconducting wire with rapid heating and quenching heat-treatment. Supercond Sci Technol 2016;29:015008. [11] Takeuchi T, Tagawa K, Kiyoshi T, Itoh K, Kosuge M, Yuyama M, et al. Enhanced current capacity of jelly-roll processed and transformed Nb3Al multifilamentary conductors. IEEE Trans Appl Supercond 1999;9:2682–7. [12] Gregory E, Tomsic M, Buta F, Sumption MD, Collings EW. Process development and microstructures of Nb3Al precursor strand for reel-to-reel production. IEEE Trans Appl Supercond 1999;9:2692–5.
4. Conclusion In this work, we systemically compared the phase transformation feature and superconducting properties of Nb3Al wire rapidly heated at 272 A and 315 A which corresponds to the current condition of getting Tc peak and valley, respectively. With the increase of annealing 6
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