Improvement in the transport critical current density and microstructure of isotopic Mg11B2 monofilament wires by optimizing the sintering temperature

Superconducting wires are widely used in fabricating magnetic coils in fusion reactors. In consideration of the stability of 11B against neutron irradiation and lower induced radio-activation properties, MgB2 superconductor with 11B serving as boron source is an alternative candidate to be used in fusion reactor with severe irradiation environment. In present work, a batch of monofilament isotopic Mg11B2 wires with amorphous 11B powder as precursor were fabricated using powder-in-tube (PIT) process at different sintering temperature, and the evolution of their microstructure and corresponding superconducting properties was systemically investigated. Accordingly, the best transport critical current density (Jc) = 2 × 104 A/cm2 was obtained at 4.2 K and 5 T, which is even comparable to multi-filament Mg11B2 isotope wires reported in other work. Surprisingly, transport Jc vanished in our wire which was heat-treated at excessively high temperature (800 °C). Combined with microstructure observation, it was found that lots of big interconnected microcracks and voids that can isolate the MgB2 grains formed in this whole sample, resulting in significant deterioration in inter-grain connectivity. The results can be a constructive guide in fabricating Mg11B2 wires to be used as magnet coils in fusion reactor systems such as ITER-type tokamak magnet.

Scientific RepoRts | 6:36660 | DOI: 10.1038/srep36660 taken into account. Compared with conventional Nb-based superconductors, MgB 2 features "low activation" and a much shorter decay time. Within 1 year, the dose rate of MgB 2 materials will be reduced to the hands-on maintenance level, which is considered as desirable for a fusion reactor magnet system 3 . Additionally, because of the reaction 10 B + n → 7Li + He (gas) under the heavy irradiation condition, 10 B can no longer guarantee the stability of the MgB 2 superconducting magnet. By replacing 10 B with the isotope 11 B, Mg 11 B 2 superconducting wires will be much more stable in a neutron irradiation environment due to the smaller neutron capture cross-section of 11 B 5 . Considering the abundant reserves of 11 B on Earth (20 wt% for 10 B, 80 wt% for 11 B), the anticipated cost for extracting the isotope from natural boron is expected to be decreased during the chemical synthesis.
The superconductivity of MgB 2 was discovered in 2001 6 . It is well-known for its simple binary chemical composition and much higher critical transition temperature (T c ) of 39 K than that of NbTi at 9.3 K. In order to operate Nb-based low-temperature superconductors, the core of the magnet needs to be cooled down to 4 K. The only eligible cryogen is liquid helium, which is extremely expensive, not always available on hand, and very difficult to handle. In the case of MgB 2 , a working temperature as high as 20 K is low enough to achieve acceptable performance. Remarkably, the operating cost is expected to be cut by over 50% by substituting cryocooler-cooled MgB 2 materials for liquid-helium-cooled Nb-based superconductors. Furthermore, the fabrication cost of MgB 2 superconducting wire itself ($2.64/kA•m) is less than 1/3 of that of Nb 3 Sn wire ($9/kA•m). Therefore, due to the advantages of cost-effectiveness, lower radioactivation, and the shorter decay time of isotopic Mg 11 B 2 , fundamental research on Mg 11 B 2 superconducting wires will be valuable for improving the efficiency of practical application in high-irradiation environments such as fusion reactors.
Mg 11 B 2 wires using isotopically pure 11 B powder always show lower critical current density (J c ) values, however, than the wires fabricated with natural boron powder. According to previous work 7,8 , this lower J c is a result of the increased amount of non-reactive precursor, which decreases the superconducting fraction. On the other hand, inter-grain connectivity is considered another crucial factor in the current-carrying capability of Mg 11 B 2 superconducting wires [9][10][11] . In this work, with the aim of further improving J c in Mg 11 B 2 wires, the evolution of the microstructure and superconducting performance in Mg 11 B 2 wires sintered at different temperatures was investigated in detail. The influence of both the superconducting fraction and the inter-grain connectivity on the J c performance is discussed. We optimized the temperature of the heat-treatment at which the best transport performance can be obtained. Surprisingly, in the case of Mg 11 B 2 wire sintered at high temperature, the transport J c vanished, although magnetic J c was still detected. According to detailed microstructure observations, this could be ascribed to the formation of a unique microstructure that was only obtained in the sample sintered at excessively high temperature. This kind of microstructure leads to significant deterioration in inter-grain connectivity and ultimately, poor transport current performance.

Experimental Details
The standard in-situ powder-in-tube (PIT) procedure was applied to all the samples. The starting materials for the Mg 11 B 2 wire consisted of 11 B amorphous powder (from Pavezyum Kimya, Turkey, Moissan method 12 , 95.5%) and magnesium powder (100-200 mesh, 99%). The isotopic purity and particle size with respect to the 11 B enriched boron powder was > 99.5% and 840 nm, respectively. After mixing the precursor powders, the mixture was tightly packed into Nb/Monel tubes with 10 mm outer diameter and 6 mm inner diameter. The composite wire was swaged and drawn to a final outer diameter of 1.08 mm. Then, the fabricated Mg 11 B 2 wires were sintered at different temperatures ranging from 700 °C, 750 °C, 770 °C, and 800 °C for 60 min (ramp rate: 5 °C/min) under high purity flowing argon gas. Finally, the samples were furnace-cooled to room temperature.
The transport critical current (I c ) measurements were carried out by using an American Magnetics superconducting magnet with DC current (with the upper limit of the current source 200 A) under possible magnetic field up to 15 T, with the standard four-probe method and the criterion of 1 μ V/cm. The critical current density J c was calculated by dividing I c by the cross-section of the Mg 11 B 2 core, which was examined with an optical microscope (Leica M205A). Scanning electron microscopy (SEM, JEOL JSM-6490LV & JEOL JSM-7500) was employed to observe the microstructure under different magnifications. X-ray diffraction (XRD) θ-2θ scans (GBC-MMA) were used to identify the phase composition. Measurements of electrical resistivity and magnetic moment were conducted in a 9 T Physical Properties Measurement System (PPMS, Quantum Design). In case of XRD, SEM, and PPMS measurements, the outer sheaths of the Mg 11 B 2 /Nb/Monel wires were removed for better data accuracy.

Results and Discussion
Typical transport J c -B performances of all four wires sintered at different temperatures are shown in Fig. 1. For reference purposes, transport J c data of for the multi-filament Mg 11 B 2 /Ta/Cu wire reported by Hishinuma 7 is also plotted in the figure. It should be noted that our best monofilament Mg 11 B 2 wire shows comparable transport J c performance to the multifilament wire fabricated by the National Institute for Fusion Science (NIFS) 7 . This result is considered as a big breakthrough, and it strongly supports the feasibility of replacing commercial NbTi by high-performance Mg 11 B 2 wires in highly radioactive fusion reactors. In our Mg 11 B 2 wires, 750 °C is the optimized temperature for heat treatment. The corresponding wire possesses a J c value near 2 × 10 4 A/cm 2 at 4.2 K and 5 T. Slight J c degradation is observed in the wire treated at temperatures deviating from 750 °C. Surprisingly, no transport current was detected in the wire treated at 800 °C. For verification, five attempts at measurement were carried out on three batches of wires produced under the same sintering conditions. Ultimately, none of them gave detectable transport current data. It is speculated that some unexpected qualitative change inside the wire might occur once the heating temperature reaches a certain level. This should probably be attributed to a unique property of the 11 B starting powder. It is believed that investigations of the phase composition, microstructure, and inter-grain connectivity will give an explanation for this abnormal phenomenon.
Scientific RepoRts | 6:36660 | DOI: 10.1038/srep36660 To confirm the phase composition, Mg 11 B 2 cores were removed from their outer sheaths and finely ground as XRD specimens. In Fig. 2(a), the main peaks indexed as Mg 11 B 2 can be observed in all spectra, indicating that the temperature is high enough to permit the formation of Mg 11 B 2 phase. Very little oxidation was detected, according to the negligible MgO peak. Un-reacted Mg and 11 B-rich phase are found in the wire sintered at relatively low temperature. Apparently, it is very hard for Mg to completely diffuse into boron particles, if the sintering temperature is not high enough. A diminishing gradient of Mg concentration exists along the radial direction of the boron particle. As a result, Mg 11 B 2 phase can only be formed on the outer layers of boron particles. The rest of the Mg will either stay in the elemental state (un-reacted Mg) or participate in other secondary reactions. Hence, 11 B-rich phase is prone to form in this case, which can be deduced from the Mg-B phase diagram 13 . The presence of those impurities (un-reacted Mg and 11 B-rich phase) will reduce the fraction of superconducting phase, which is crucial for the final performance of superconductors. It has to be pointed out that the chemical activity of 11 B is lower in comparison with natural B due to the isotope kinetic effect 14 . This might explain why 700 °C is not high enough for the complete reaction in this work. Figure 2(b) shows the mass fractions of Mg 11 B 2 phase in the wires as a function of sintering temperature. The mass fractions were calculated by using Rietveld refinement. The smallest Mg 11 B 2 fraction, as low as 84.7%, is found in the wire treated at 700 °C. This is mainly due to the presence of impurities, as reflected by the XRD results. Furthermore, the degradation in transport J c performance also confirms its relatively poor superconductivity (see Fig. 1). With increasing sintering temperature, un-reacted Mg peaks become smaller and almost disappear. Correspondingly, the mass fractions of Mg 11 B 2 phase in the rest of the wires all remain at a high level (> 90%). Since the crystallization of Mg 11 B 2 phase is confirmed to be good in the Mg 11 B 2 wire sintered at 800 °C, while its mass fraction of superconducting phase is also satisfactory, the observed abrupt disappearance of transport current in the wire sintered at 800 °C is related to neither the phase composition nor a low superconducting fraction.    , and a and b (a < b) are the length and width of the cross-section which is perpendicular to the direction of the applied magnetic field 15 . In our case, the Mg 11 B 2 cores are cylindrical in shape. So, the formula can be simplified to J c = 30(Δ M/V)/d, where d is the diameter of the circular cross-sectional area 16,17 . According to the calculations, the magnetic J c (H) results at 5.0 K are shown in Fig. 3(b). The wire sintered at 750 °C shows the best magnetic J c (H) performance throughout the entire range of fields, which is consistent with the transport J c results shown in Fig. 1. Some differences can be found between the values of magnetic J c and transport J c . Other than measurement deviation, the intrinsic distinction between the magnetic J c signal and the transport J c signal also needs to be taken into consideration. Generally, due to the existence of negative structures such as porosity and cracks, not all the MgB 2 in a sample is capable of passing transport current. Inter-or intra-grain connectivity should always be considered when dealing with transport performance. On the contrary, as long as they possess superconductivity, all the MgB 2 fragments will contribute to the magnetic J c . It should be noted that the magnetic J c was detected and showed good performance in the wire sintered at 800 °C. This means that the Mg 11 B 2 superconducting phase in the wire was not badly damaged by the high sintering temperature. Hence, after ruling out the effects of inferior superconducting phase, it can be speculated that the transport current in the wire sintered at 800 °C disappeared as a result of a problem with inter-grain connectivity. A high sintering temperature might introduce some defects and significantly destroy the connection between Mg 11 B 2 superconducting grains.
It is estimated that the vanishing of transport current in the Mg 11 B 2 wire is caused by the severe deterioration of inter-grain connectivity, which can be visually confirmed by SEM micrographs. The low-magnification SEM images of the cross-sections of Mg 11 B 2 wires sintered at 700 °C, 750 °C and 800 °C are presented in Fig. 4(a-c). Obvious evolution of the surface morphology is exhibited with increasing temperature. In the wire treated at 700 °C, it was already proved by the XRD results that the Mg had partially reacted with the boron. As the particle size of the Mg powder is much bigger than for the boron powder, un-reacted Mg melted and smoothly covered the Mg 11 B 2 grains. Therefore, the morphology of this sample was fairly plain and incompact. A dense surface is observed in Fig. 4(b) on the optimal sample sintered at 750 °C, indicating complete reaction and good inter-grain connectivity. This is consistent with the J c -B and XRD results discussed above. Once the sintering temperature reached 800 °C, big cracks (marked by black arrows) were observed, as shown in Fig. 4(c). They are much bigger than the normal microcracks in other samples. Note that most of the big cracks are connected with each other. This feature is considered to be highly detrimental to the inter-grain connectivity. The resultant superconducting fragments are isolated from each other, and eventually, very little current can pass through the entire wire, which will significantly reduce the transport performance. On further increasing the magnification, porous structure is found in the same sample (marked by white arrows in Fig. 4(d)). When the wire was heat-treated at 800 °C, both the grain size and the mobility of the Mg 11 B 2 grains were increased. The separate grains are prone to aggregate with each other, leaving plenty of voids in the morphology. Consequently, the effective current capacity is sharply reduced with the emergence of the porous structure. This is considered to be another barrier to obtaining high transport current in Mg 11 B 2 wires. In addition, this kind of microstructure with abundant voids can be more brittle and thus be more prone to fracture and form big microcracks (see Fig. 4c) resulting from heat stress during the furnace-cooling process from high temperature to room temperature.
High-resolution SEM was employed to investigate the details of the crystalline structure in the four Mg 11 B 2 , wires, and the results are presented in Fig. 5. In the sample with the lowest sintering temperature, the crystalline grains have a wide range of sizes, and all of them are dispersed in a melted matrix, as shown in Fig. 5(a). Referring to the XRD results above, the melted matrix is un-reacted Mg, which cannot be fully reacted with B at a relatively low temperature. This is strong evidence for the smaller mass fraction of Mg 11 B 2 phase and lower transport performance in this sample. In the wires sintered at higher temperature, the amount of un-reacted Mg is greatly reduced, and the Mg 11 B 2 crystalline grains keep growing and form typical hexagonal shapes, which can be observed in Fig. 5(b,c). Figure 5(d) shows the morphology of the wire sintered at 800 °C, in which some grains abnormally grow, and abundant big clusters are found. These clusters are formed by the localized aggregation of Mg 11 B 2 grains at the relatively high heat-treatment temperature. This phenomenon further increases the porosity on the macroscale and significantly reduces the effective superconducting fraction for transporting current. As a result, the inter-grain connectivity is badly degraded. Combining these results with the low-magnification SEM images, it is thus concluded that the vanishing of transport current in the Mg 11 B 2 sintered at high temperature should be attributed to the depression of inter-grain connectivity in the wire that is caused by the big microcracks and high porosity.

Conclusions
The effects of sintering temperature on the superconducting performance and morphology of Mg 11 B 2 monofilament wires made from isotopically pure boron powder were investigated in this work. It was found that increasing the sintering temperature led to the evolution of microstructure and characteristic changes in the transport current capacity. Un-reacted Mg and B-rich phase existed in the wire sintered at low temperature. The Mg 11 B 2 fraction, as well as the transport performance, was reduced because of the un-reacted Mg and B-rich phase impurities. With increasing sintering temperature, better phase composition and crystallinity were obtained. The best transport J c = 2 × 10 4 A/cm 2 was reached at 4.2 K and 5 T in the Mg 11 B 2 wire sintered at 750 °C. It should be noted that although high magnetic J c was detected in the wire sintered at 800 °C, the transport current was totally absent. The evolution of the morphology could be clearly seen in the wires corresponding to different sintering temperatures. Due to the abnormal growth and high mobility of Mg 11 B 2 grains at relatively high ambient temperature, numerous big microcracks, voids, and Mg 11 B 2 clusters formed in the wire sintered at 800 °C. As a result, the inter-grain connectivity was significantly suppressed, resulting in the inferior transport performance. The results obtained in our work can be a constructive guide for fabricating Mg 11 B 2 wires to be used as magnet coils in fusion reactor systems such as ITER-type tokamak magnets.