- Email: [email protected]
Properties of MgB2 wires made of oxidized powders
Physica C 477 (2012) 20–23
Contents lists available at SciVerse ScienceDirect
Physica C journal homepage: www.elsevier.com/locate/physc
Properties of MgB2 wires made of oxidized powders P. Kovácˇ a,⇑, M. Kulich a, W. Haessler b, M. Hermann b, T. Melišek a, M. Reissner c a
Institute of Electrical Engineering, Slovak Academy of Sciences, Dúbravska cesta 9, 841 04 Bratislava, Slovakia Institute for Metallic Materials, Leibniz-Institute for Solid State and Materials Research (IFW), Dresden, Germany c Institute of Solid State Physics, Vienna University of Technology, Vienna, Austria b
a r t i c l e
i n f o
Article history: Received 16 December 2011 Received in revised form 1 February 2012 Accepted 20 February 2012 Available online 1 March 2012 Keywords: MgB2 In situ process Oxidized powders MgO Critical currents Resistivity
a b s t r a c t Single-core MgB2/Ti wires have been made from mixed (in situ) and intensively milled (mechanically alloyed) Mg + B powders. Both powders were oxidized at room temperature for 48 h and also by heat treatment at 400 °C/1 h in air. Critical current density and resistivity measurements were performed for the samples with oxidized powders and also for those protected in glove box with high purity Ar. It was found that oxidation of in situ mixture is slightly decreasing critical current density. Surprisingly, the oxidation of much finer mechanically alloyed powder is improving Jc in spite of increased MgB2 resistivity. Jc at 10T increase by 47–100% was measured for mechanically alloyed powder oxidized at room temperature. This allows a safety handling with intensively milled Mg + B powders in air. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction It is well known that oxidation of Mg is very easy and starting at low temperatures [1]. Therefore, it is impossible to prepare MgB2 phase without any MgO. Magnesium oxide is often present as an impurity in commercial MgB2 powder and also formed during heat treatment of wires made by ex situ and in situ process [2–5]. MgO is formed usually by the reaction of Mg vapour (e.g. from decomposed MgB2 at high temperature) with O2 gas (trapped in pores during processing, or contained in the annealing gas) and can be present in crystalline and/or amorphous phase. Volume of trapped oxygen is strongly affected by the powder grain size, which is in inverse relation to the surface area. X-ray diffraction is frequently used to detect MgO impurity levels, but the detection limit for crystalline phase is around 2–3% and amorphous one is not detectable [6]. In the case of milling, the surface area of ex situ and in situ powders is increasing with milling energy and time [7–8], which may influence the content of MgO during the next steps. Attempts to remove MgO from MgB2 powder by chemical etching with benzoic and acetic acid have been done [9–11]. It was shown that high MgO content prevents the current flow, which is attributed to its location in the grain boundaries [12]. Therefore, careful handling with powders is usually used for high performance MgB2 materials. Vignolo et al. have studied the effect of oxidizing atmosphere in ex situ tapes [13]. They have found that MgO concentration in MgB2 ⇑ Corresponding author. Tel.: +421 2 5922 2841; fax: +421 2 5477 5816. E-mail address: [email protected] (P. Kovácˇ). 0921-4534/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.physc.2012.02.026
phase has a great importance and yields a positive or negative contribute on behaviour in low and high magnetic fields. Up to now, most of oxidation studies were done for ex situ process. The aim of this contribution is to present the effect of oxidation for in situ and mechanically alloyed powders. 2. Experimental Single-core Ti sheathed wires have been made using in situ (I) (mixed – 5 h at 100 rpm in the ball mill, powder to ball ratio 1:18), mechanically alloyed (M) (high energy milled – 50 h, 250 rpm, powder to ball ratio 1:36) and ex situ (E) powders (described by Table 1). Milling and mixing was done in tungstencarbide tools under inert Ar atmosphere in a planetary ball mill (Retsch PM4000) [14]. Reference samples were made by Ti-tube (6/4 mm) filling with powder kept in glove box with high purity Argon atmosphere (called I, M and E wires). The same powders (Mg from Goodfellow <250 lm and Boron from Fluka (95–97% purity) were used for wires made of oxidized powders: at room temperature for 48 h (Irt48 and Mrt48) and also heat treated in air at 400 °C/1 h (Iht1 and Mht1). All wires were deformed by rotary swaging to 3.0 mm and then by two-axial rolling up to 1.0 1.0 mm2. Wires made of I and M powders were finally annealed at 625 °C/3 h or 650 °C/0.5 h and wire E at 950 °C/0.5 h. Critical current densities (Jc) were measured at 4.2 K in the range of external fields 4–10T. Temperature dependences of MgB2 core resistivity (extracted from Ti sheath) was measured between 30 and 300 K. Critical temperature (Tc) was defined as the middle of
P. Kovácˇ et al. / Physica C 477 (2012) 20–23
21
Fig. 1a shows the critical current density dependences for three wires differing by used powder (I, M and E), which were protected in glove box with high purity Argon atmosphere. E wire exhibits the softest in-field Jc dependence with interesting values only in low external fields, B < 3T. Less steeped Jc(B) is measured for I wire
but Jc magnitudes are the lowest for B < 7.5T, which is attributed to higher porosity (lower density) typical for a standard in situ process [15]. The best Jc(B) performance is evident for M wire due to highly dense and fine-grain structure obtained by the deformation of mechanically alloyed powder [16]. Fig. 1b shows the Kramer’s plot for compared wires with an extrapolation to Birr giving: 10.4, 11.3 and 12.3T for E, M and I wire, respectively. Superconducting parameters reflect the differences in grain-connectivity and pinning affected by the core density, grain size and impurities (MgO, Mg,. . .) distribution and location. Very detailed and demanding structural analysis would be needed for a complete understanding of presented differences. Fig. 2a shows Jc(B) characteristics of I and M wires with oxidized powders heat treated at 625 °C/3 h. While in the case of I wire both oxidations have reduced slightly the current density, increased Jc were measured for M wires oxidized by heat treatment at 400 ° C/1 h and most effectively by the oxidation at room temperature for 48 h (see inset in Fig. 2a). More expressive changes in Jc(B) have been observed for the heat treatment performed above the melting point of magnesium (650 °C/0.5 h), which is presented by Fig. 2b. Higher Jc reduction for Irt48 and Iht1 and larger Jc increase for Mrt48 and Mht1 is apparent (see inset in Fig. 2b). This is surprising result because the surface area of M powder is much larger than I and therefore applied oxidations should increase the MgO content more for M. Consequently, reduced Jc were expected for Mrt48 and Mht1 wires. More than one order of magnitude Jc degradation
Fig. 1. Critical current density for wire samples differing only by used powder (I, M and E) which were protected in glove box with high purity Argon atmosphere (a). I and M wires were heat treated at 625 °C/3 h and the wire E at 850 °C/0.5 h. Kramer’s plot made for I, M and E wires (b).
Fig. 2. Jc(B) characteristics of I and M wires with oxidized powders heat treated at 625 °C/3 h (a) and at 650 °C/0.5 h (b). The insets show the effect of applied oxidation on Jc(10T).
Table 1 Description of powders used for Ti sheathed MgB2 wires. I Irt48 Iht1 M Mrt48 Mht1 E
Mixture of Mg + B, kept in glove box with high purity Ar I – mixture, oxidized at room temperature for 48 h I – mixture, air heated at 400 °C/1 h I – mixture milled at 250 rpm for 50 h, kept in glove box with high purity Ar M-powder, oxidized at room temperature for 48 h M-powder, air heated at 400 °C/1 h Commercial MgB2 Alfa Aesar powder, kept in glove box with high purity Ar
resistive transition and DTc as the width of transition. Phase composition of powdered MgB2 cores was analysed by X-ray and by following Rietveld refinement [14].
3. Results and discussion 3.1. Critical current densities
22
P. Kovácˇ et al. / Physica C 477 (2012) 20–23
was observed for MgB2 wire made of ex situ powder oxidized by heat treatment in air at 500 °C/1 h [11]. Such high reduction of current density is attributed to the oxidation of grain surfaces protecting the inter-grain current flow. In the case of I and M wires, oxidized Mg particles are co-deformed with boron and finally converted into MgB2 phase by heat treatment. Due to different particle size of initial powder mixtures (Mg < 250 lm + < 1 lm B for I and nano-size Mg + B + MgB2 (30%) for M [14]), the content and morphology of created MgO should be unequal, which influences the transport currents of I and M wires differently. Small (nano-size) and uniformly distributed MgO particles can possibly act as effective pinning centres increasing in-field Jc [3], which may be a possible reason of observed in-field Jc improvement for Mrt48 and Mht1 wires at both heat treatment temperatures. No in-field Jc improvement has been observed for Irt48 and Iht1 wires. Jiang et al. have presented Jc improvement in in situ MgB2 by nano-size MgO doping with optimal content 1.25 wt.% for in situ wire [17]. They reached Jc(10T) enhancement by 50% which is comparable to the presented one showing 47% improvement of Jc(10T) for wire Mrt48 heat treated at 625 °C/3 h (see inset in Fig. 2a). Higher increase of Jc(10T) (by 100%) is measured for Mrt48 wire heat treatment at 650 °C/0.5 h (see inset in Fig. 2b). It indicates that liquid-phase reaction may introduce more MgO particles inside the MgB2 grains and less at grain boundaries. But this hypothesis is not yet verified by structural (TEM) observations. The lowest Jc(10T) values were measured for Iht1, which can be a reason of the largest amount of MgO located at the grain boundaries.
3.2. Resistivities of MgB2 cores To analyse the effect of oxidation more in details, MgB2 cores have been extracted from Ti sheath and the resistively measured by four-probe arrangement [18]. Fig. 3 shows the temperature dependent qMgB2 for oxidized cores compared to the reference ones (I and M). It is apparent that both oxidations have increased the resistivity qMgB2. While the q(T) dependences of M wires shows a typical shape measured usually for MgB2 samples [19], all I wires have nearly linear decrease of resistivity with temperature up to the point of superconducting transition. This indicates different phase compositions in comparison of M with I wires. Table 2 presents critical temperatures, Tc, and resistivity data for compared MgB2 cores. One can see that Tc of M wires is slightly decreased by both powder oxidations. Nearly unchanged Tc’s were measured for wires Iht1 and Irt48, but the width of superconducting transition, DTc, is considerably larger for I wires indicating less pure MgB2. Rowell analysis has been applied for all presented
Table 2 Critical temperature and resistivity data for M and I samples annealed at 625 °C/3 h. Sample
Tc (K)
DTc (K)
Dq (lX cm)
AF
RRR
M Mht1 Mrt48 I Iht1 Irt48
36.8 36.5 36.6 37.6 37.6 37.7
0.2 0.3 0.2 0.6 0.7 0.8
152.4 166.6 176.9 284.6 316.2 333.1
0.048 0.044 0.041 0.026 0.023 0.022
1.90 1.82 1.85 8.79 5.00 4.73
samples [19]. The measured resistivity values at 40 and 300 K were used for phonon term resistivity Dq, which is the change in the q values at room temperature and just above the Tc, i.e., at 40 K. The fraction of effective current carrying cross-sectional area, AF, or connectivity, is estimated by comparing phonon term resistivity Dq of a perfectly connected ideal specimen (7.3 lO cm [20]) with that of the observed value of the sample. Dq, AF, and the residual resistivity ratio RRR, are all gathered in Table 2. It is evident that M and I wires are differing by all estimated parameters. Values of AF, varying from 0.022 to 0.026 for I samples due to poor grain connectivity likely attributed to the presence of pores and insulating phase MgO. The general inspection of AF reveals that the used oxidation reduces the connectivity between grains. AF values for M samples are always larger (0.041–0.048) than those of the equivalent I samples. This can partly be due to denser (less porous) MgB2 core in M wires. Fig. 2a presents Jc data not correlating with the q and AF (increased q ? improved Jc). These results indicate, contrary to what was proposed by Rowell, that the q or AF values cannot be directly used to predict the Jc(B) performance. Much higher RRR measured for I wire are due to different phase composition, which is discussed Section 3.3.
3.3. Phase composition of MgB2 cores The basic structural parameters and the content of phases for selected samples acquired using the X-ray and Rietveld refinement is presented in Tables 3a and 3b. None of the analysed cores contain detectable amounts of MgB4 but the amount of MgO varies from 4% to 12% and un-reacted Mg from 5% to 17% depending on the used powder and applied oxidation. Generally, X-ray analysis shows a comparable volume of crystalline MgO for I and M samples although the surface area of these two powders are different. Increased content of MgO (doubled) is evident for all oxidized samples. On the other side, higher content of MgB2 phase is found for M wires and large amount of not reacted Mg (12–17%) in I wires, which is responsible for not typical q(T) dependences shown in Fig. 3 and also for much higher RRR in Table 2. Due to high content of not reacted Mg in Irt48 and Iht1 samples, the connectivity estimated by Rowell approach analysis used for all presented samples are probably not fully correct. Instead of that, the parameters proposed by Rowell (Dq or AF) are not correlating well with Jc data and cannot be directly used to predict the Jc(B) performance, which is influenced not only by the connectivity
Table 3a Phase composition from X-ray and Rietveld refinement for M and I samples annealed at 625 °C/3 h.
Fig. 3. The resistivities of MgB2 cores from oxidized powders compared to the reference ones.
Sample
MgB2 (%)
MgO (%)
Mg (%)
M Mht1 Mrt48 I Iht1 Irt48
88.0 84.0 80.0 88.0 80.0 78.0
4.0 11 12 5.0 8.0 9.0
8.0 5.0 8.0 17 12 13
P. Kovácˇ et al. / Physica C 477 (2012) 20–23 Table 3b Phase composition of M and I samples annealed at 650 °C/0.5 h.
23
Acknowledgement
Sample
MgB2 (%)
MgO (%)
Mg (%)
M Mht1 Mrt48 I Iht1 Irt48
85.0 83.0 84.0 78.0 78.0 76.0
4.0 10.0 7.0 5.0 10.0 8.0
11.0 7.0 9.0 17.0 12.0 16.0
but also by pinning. Positive role of small and well distributed MgO particles is considered as a main reason of increase Jc in Mrt48 and Mht1 wires.
4. Conclusions Basic properties of Ti sheathed MgB2 wires made of oxidized in situ and mechanically alloyed powders have been studied. It was found that oxidation of Mg + B mixture in air is not as risky as for ex situ made MgB2 and decreasing the critical current density only slightly. Surprisingly, the oxidation of intensively milled (mechanically alloyed) powder is improving critical current density especially in high magnetic field (Jc at 10T increased by 47–100%). It enables a safety handling with fine grain mechanically alloyed powder in air. It is also shown that widely used Rowell analysis (for grain connectivity) cannot be used generally for prediction of Jc(B) performance of any kind of MgB2 samples, which are influenced also by the pinning at grain boundaries and at nanosize and well distributed defects.
This work was supported by the Slovak Scientific Agency under the Project APVV-0495-10. References [1] M. Kurth, et al.. [2] Y. Zhu, L. Wum, V. Volkov, Q. Li, G. Gu, A.R. Moodenbaugh, M. Malac, M. Suenaga, J. Tranq. Physica C 356 (2001) 239. [3] P. Kovácˇ, I. Hušek, T. Melišek, J.C. Grivel, W. Pachla, V. Štrbík, R. Diduszko, J. Homeyer, N.H. Andersen, Sup. Sci. Technol. 17 (2004) L41. [4] O. Perner, J. Eckert, W. Häßler, C. Fischer, K.-H. Müller, G. Fuchs, B. Holzapfel, L. Schultz, Supercond. Sci. Technol. 17 (2004) 1148. [5] L. Gu et al., J. Cryst. Growth 280 (2005) 602. [6] C.E.J. Dancer, P. Mikheenko, A. Bevan, J.S. Abell, R.I. Todd, C.R.M. Grovenor, J. Eur. Ceram. Soc. 29 (2009) 1817. [7] A. Kario, R. Nast, W. Häßler, C. Rodig, C. Mickel, W. Goldacker, B. Holzapfel, L. Schultz, Supercond. Sci. Technol. 24 (2011) 075011. [8] B.J. Senkowicz et al., Supercond. Sci. Technol. 21 (2008) 035009. [9] H. Fujii, K. Togano, K. Ozawa, J. Mater. Res. 22 (2007) 1281. [10] K. Hušekova, I. Hušek, P. Kovácˇ, M. Kulich, E. Dobrocˇka, V. Štrbík, Physica C 470 (2010) 331. [11] P. Kovácˇ, I. Hušek, M. Kulich, T. Melišek, K. Hušeková, E. Dobrocˇka, Physica C 470 (2010) 340. [12] J. Jiang et al., Supercond. Sci. Technol. 19 (2006) L33. [13] M. Vignolo et al., IEEE Trans. Appl. Supercond. 19 (2009) 2718. [14] M. Herrmann, W. Häßler, C. Mickel, W. Gruner, B. Holzapfel, L. Schultz, Supercond. Sci. Technol. 20 (2007) 1108. [15] A.V. Pan, S. Zhou, H. Liu, S.X. Dou, Supercond. Sci. Technol. 16 (2003) 639. [16] W. Häßler, M. Herrmann, C. Rodig, M. Schubert, K. Nenkov, B. Holzapfel, Supercond. Sci. Technol. 21 (2008) 062001. [17] C.H. Jiang, H. Hatakeyama, H. Kumakura, Physica C 423 (2005) 45. [18] M. Kulich, P. Kovácˇ, M. Eisterer, I. Hušek, H. Weber, W. Haessler, Physica C 469 (2009) 827. [19] J.M. Rowell, Supercond. Sci. Technol. 16 (2003) R17. [20] R.H.T. Wilke et al., Physica C 424 (2005) 1.
Contents lists available at SciVerse ScienceDirect
Physica C journal homepage: www.elsevier.com/locate/physc
Properties of MgB2 wires made of oxidized powders P. Kovácˇ a,⇑, M. Kulich a, W. Haessler b, M. Hermann b, T. Melišek a, M. Reissner c a
Institute of Electrical Engineering, Slovak Academy of Sciences, Dúbravska cesta 9, 841 04 Bratislava, Slovakia Institute for Metallic Materials, Leibniz-Institute for Solid State and Materials Research (IFW), Dresden, Germany c Institute of Solid State Physics, Vienna University of Technology, Vienna, Austria b
a r t i c l e
i n f o
Article history: Received 16 December 2011 Received in revised form 1 February 2012 Accepted 20 February 2012 Available online 1 March 2012 Keywords: MgB2 In situ process Oxidized powders MgO Critical currents Resistivity
a b s t r a c t Single-core MgB2/Ti wires have been made from mixed (in situ) and intensively milled (mechanically alloyed) Mg + B powders. Both powders were oxidized at room temperature for 48 h and also by heat treatment at 400 °C/1 h in air. Critical current density and resistivity measurements were performed for the samples with oxidized powders and also for those protected in glove box with high purity Ar. It was found that oxidation of in situ mixture is slightly decreasing critical current density. Surprisingly, the oxidation of much finer mechanically alloyed powder is improving Jc in spite of increased MgB2 resistivity. Jc at 10T increase by 47–100% was measured for mechanically alloyed powder oxidized at room temperature. This allows a safety handling with intensively milled Mg + B powders in air. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction It is well known that oxidation of Mg is very easy and starting at low temperatures [1]. Therefore, it is impossible to prepare MgB2 phase without any MgO. Magnesium oxide is often present as an impurity in commercial MgB2 powder and also formed during heat treatment of wires made by ex situ and in situ process [2–5]. MgO is formed usually by the reaction of Mg vapour (e.g. from decomposed MgB2 at high temperature) with O2 gas (trapped in pores during processing, or contained in the annealing gas) and can be present in crystalline and/or amorphous phase. Volume of trapped oxygen is strongly affected by the powder grain size, which is in inverse relation to the surface area. X-ray diffraction is frequently used to detect MgO impurity levels, but the detection limit for crystalline phase is around 2–3% and amorphous one is not detectable [6]. In the case of milling, the surface area of ex situ and in situ powders is increasing with milling energy and time [7–8], which may influence the content of MgO during the next steps. Attempts to remove MgO from MgB2 powder by chemical etching with benzoic and acetic acid have been done [9–11]. It was shown that high MgO content prevents the current flow, which is attributed to its location in the grain boundaries [12]. Therefore, careful handling with powders is usually used for high performance MgB2 materials. Vignolo et al. have studied the effect of oxidizing atmosphere in ex situ tapes [13]. They have found that MgO concentration in MgB2 ⇑ Corresponding author. Tel.: +421 2 5922 2841; fax: +421 2 5477 5816. E-mail address: [email protected] (P. Kovácˇ). 0921-4534/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.physc.2012.02.026
phase has a great importance and yields a positive or negative contribute on behaviour in low and high magnetic fields. Up to now, most of oxidation studies were done for ex situ process. The aim of this contribution is to present the effect of oxidation for in situ and mechanically alloyed powders. 2. Experimental Single-core Ti sheathed wires have been made using in situ (I) (mixed – 5 h at 100 rpm in the ball mill, powder to ball ratio 1:18), mechanically alloyed (M) (high energy milled – 50 h, 250 rpm, powder to ball ratio 1:36) and ex situ (E) powders (described by Table 1). Milling and mixing was done in tungstencarbide tools under inert Ar atmosphere in a planetary ball mill (Retsch PM4000) [14]. Reference samples were made by Ti-tube (6/4 mm) filling with powder kept in glove box with high purity Argon atmosphere (called I, M and E wires). The same powders (Mg from Goodfellow <250 lm and Boron from Fluka (95–97% purity) were used for wires made of oxidized powders: at room temperature for 48 h (Irt48 and Mrt48) and also heat treated in air at 400 °C/1 h (Iht1 and Mht1). All wires were deformed by rotary swaging to 3.0 mm and then by two-axial rolling up to 1.0 1.0 mm2. Wires made of I and M powders were finally annealed at 625 °C/3 h or 650 °C/0.5 h and wire E at 950 °C/0.5 h. Critical current densities (Jc) were measured at 4.2 K in the range of external fields 4–10T. Temperature dependences of MgB2 core resistivity (extracted from Ti sheath) was measured between 30 and 300 K. Critical temperature (Tc) was defined as the middle of
P. Kovácˇ et al. / Physica C 477 (2012) 20–23
21
Fig. 1a shows the critical current density dependences for three wires differing by used powder (I, M and E), which were protected in glove box with high purity Argon atmosphere. E wire exhibits the softest in-field Jc dependence with interesting values only in low external fields, B < 3T. Less steeped Jc(B) is measured for I wire
but Jc magnitudes are the lowest for B < 7.5T, which is attributed to higher porosity (lower density) typical for a standard in situ process [15]. The best Jc(B) performance is evident for M wire due to highly dense and fine-grain structure obtained by the deformation of mechanically alloyed powder [16]. Fig. 1b shows the Kramer’s plot for compared wires with an extrapolation to Birr giving: 10.4, 11.3 and 12.3T for E, M and I wire, respectively. Superconducting parameters reflect the differences in grain-connectivity and pinning affected by the core density, grain size and impurities (MgO, Mg,. . .) distribution and location. Very detailed and demanding structural analysis would be needed for a complete understanding of presented differences. Fig. 2a shows Jc(B) characteristics of I and M wires with oxidized powders heat treated at 625 °C/3 h. While in the case of I wire both oxidations have reduced slightly the current density, increased Jc were measured for M wires oxidized by heat treatment at 400 ° C/1 h and most effectively by the oxidation at room temperature for 48 h (see inset in Fig. 2a). More expressive changes in Jc(B) have been observed for the heat treatment performed above the melting point of magnesium (650 °C/0.5 h), which is presented by Fig. 2b. Higher Jc reduction for Irt48 and Iht1 and larger Jc increase for Mrt48 and Mht1 is apparent (see inset in Fig. 2b). This is surprising result because the surface area of M powder is much larger than I and therefore applied oxidations should increase the MgO content more for M. Consequently, reduced Jc were expected for Mrt48 and Mht1 wires. More than one order of magnitude Jc degradation
Fig. 1. Critical current density for wire samples differing only by used powder (I, M and E) which were protected in glove box with high purity Argon atmosphere (a). I and M wires were heat treated at 625 °C/3 h and the wire E at 850 °C/0.5 h. Kramer’s plot made for I, M and E wires (b).
Fig. 2. Jc(B) characteristics of I and M wires with oxidized powders heat treated at 625 °C/3 h (a) and at 650 °C/0.5 h (b). The insets show the effect of applied oxidation on Jc(10T).
Table 1 Description of powders used for Ti sheathed MgB2 wires. I Irt48 Iht1 M Mrt48 Mht1 E
Mixture of Mg + B, kept in glove box with high purity Ar I – mixture, oxidized at room temperature for 48 h I – mixture, air heated at 400 °C/1 h I – mixture milled at 250 rpm for 50 h, kept in glove box with high purity Ar M-powder, oxidized at room temperature for 48 h M-powder, air heated at 400 °C/1 h Commercial MgB2 Alfa Aesar powder, kept in glove box with high purity Ar
resistive transition and DTc as the width of transition. Phase composition of powdered MgB2 cores was analysed by X-ray and by following Rietveld refinement [14].
3. Results and discussion 3.1. Critical current densities
22
P. Kovácˇ et al. / Physica C 477 (2012) 20–23
was observed for MgB2 wire made of ex situ powder oxidized by heat treatment in air at 500 °C/1 h [11]. Such high reduction of current density is attributed to the oxidation of grain surfaces protecting the inter-grain current flow. In the case of I and M wires, oxidized Mg particles are co-deformed with boron and finally converted into MgB2 phase by heat treatment. Due to different particle size of initial powder mixtures (Mg < 250 lm + < 1 lm B for I and nano-size Mg + B + MgB2 (30%) for M [14]), the content and morphology of created MgO should be unequal, which influences the transport currents of I and M wires differently. Small (nano-size) and uniformly distributed MgO particles can possibly act as effective pinning centres increasing in-field Jc [3], which may be a possible reason of observed in-field Jc improvement for Mrt48 and Mht1 wires at both heat treatment temperatures. No in-field Jc improvement has been observed for Irt48 and Iht1 wires. Jiang et al. have presented Jc improvement in in situ MgB2 by nano-size MgO doping with optimal content 1.25 wt.% for in situ wire [17]. They reached Jc(10T) enhancement by 50% which is comparable to the presented one showing 47% improvement of Jc(10T) for wire Mrt48 heat treated at 625 °C/3 h (see inset in Fig. 2a). Higher increase of Jc(10T) (by 100%) is measured for Mrt48 wire heat treatment at 650 °C/0.5 h (see inset in Fig. 2b). It indicates that liquid-phase reaction may introduce more MgO particles inside the MgB2 grains and less at grain boundaries. But this hypothesis is not yet verified by structural (TEM) observations. The lowest Jc(10T) values were measured for Iht1, which can be a reason of the largest amount of MgO located at the grain boundaries.
3.2. Resistivities of MgB2 cores To analyse the effect of oxidation more in details, MgB2 cores have been extracted from Ti sheath and the resistively measured by four-probe arrangement [18]. Fig. 3 shows the temperature dependent qMgB2 for oxidized cores compared to the reference ones (I and M). It is apparent that both oxidations have increased the resistivity qMgB2. While the q(T) dependences of M wires shows a typical shape measured usually for MgB2 samples [19], all I wires have nearly linear decrease of resistivity with temperature up to the point of superconducting transition. This indicates different phase compositions in comparison of M with I wires. Table 2 presents critical temperatures, Tc, and resistivity data for compared MgB2 cores. One can see that Tc of M wires is slightly decreased by both powder oxidations. Nearly unchanged Tc’s were measured for wires Iht1 and Irt48, but the width of superconducting transition, DTc, is considerably larger for I wires indicating less pure MgB2. Rowell analysis has been applied for all presented
Table 2 Critical temperature and resistivity data for M and I samples annealed at 625 °C/3 h. Sample
Tc (K)
DTc (K)
Dq (lX cm)
AF
RRR
M Mht1 Mrt48 I Iht1 Irt48
36.8 36.5 36.6 37.6 37.6 37.7
0.2 0.3 0.2 0.6 0.7 0.8
152.4 166.6 176.9 284.6 316.2 333.1
0.048 0.044 0.041 0.026 0.023 0.022
1.90 1.82 1.85 8.79 5.00 4.73
samples [19]. The measured resistivity values at 40 and 300 K were used for phonon term resistivity Dq, which is the change in the q values at room temperature and just above the Tc, i.e., at 40 K. The fraction of effective current carrying cross-sectional area, AF, or connectivity, is estimated by comparing phonon term resistivity Dq of a perfectly connected ideal specimen (7.3 lO cm [20]) with that of the observed value of the sample. Dq, AF, and the residual resistivity ratio RRR, are all gathered in Table 2. It is evident that M and I wires are differing by all estimated parameters. Values of AF, varying from 0.022 to 0.026 for I samples due to poor grain connectivity likely attributed to the presence of pores and insulating phase MgO. The general inspection of AF reveals that the used oxidation reduces the connectivity between grains. AF values for M samples are always larger (0.041–0.048) than those of the equivalent I samples. This can partly be due to denser (less porous) MgB2 core in M wires. Fig. 2a presents Jc data not correlating with the q and AF (increased q ? improved Jc). These results indicate, contrary to what was proposed by Rowell, that the q or AF values cannot be directly used to predict the Jc(B) performance. Much higher RRR measured for I wire are due to different phase composition, which is discussed Section 3.3.
3.3. Phase composition of MgB2 cores The basic structural parameters and the content of phases for selected samples acquired using the X-ray and Rietveld refinement is presented in Tables 3a and 3b. None of the analysed cores contain detectable amounts of MgB4 but the amount of MgO varies from 4% to 12% and un-reacted Mg from 5% to 17% depending on the used powder and applied oxidation. Generally, X-ray analysis shows a comparable volume of crystalline MgO for I and M samples although the surface area of these two powders are different. Increased content of MgO (doubled) is evident for all oxidized samples. On the other side, higher content of MgB2 phase is found for M wires and large amount of not reacted Mg (12–17%) in I wires, which is responsible for not typical q(T) dependences shown in Fig. 3 and also for much higher RRR in Table 2. Due to high content of not reacted Mg in Irt48 and Iht1 samples, the connectivity estimated by Rowell approach analysis used for all presented samples are probably not fully correct. Instead of that, the parameters proposed by Rowell (Dq or AF) are not correlating well with Jc data and cannot be directly used to predict the Jc(B) performance, which is influenced not only by the connectivity
Table 3a Phase composition from X-ray and Rietveld refinement for M and I samples annealed at 625 °C/3 h.
Fig. 3. The resistivities of MgB2 cores from oxidized powders compared to the reference ones.
Sample
MgB2 (%)
MgO (%)
Mg (%)
M Mht1 Mrt48 I Iht1 Irt48
88.0 84.0 80.0 88.0 80.0 78.0
4.0 11 12 5.0 8.0 9.0
8.0 5.0 8.0 17 12 13
P. Kovácˇ et al. / Physica C 477 (2012) 20–23 Table 3b Phase composition of M and I samples annealed at 650 °C/0.5 h.
23
Acknowledgement
Sample
MgB2 (%)
MgO (%)
Mg (%)
M Mht1 Mrt48 I Iht1 Irt48
85.0 83.0 84.0 78.0 78.0 76.0
4.0 10.0 7.0 5.0 10.0 8.0
11.0 7.0 9.0 17.0 12.0 16.0
but also by pinning. Positive role of small and well distributed MgO particles is considered as a main reason of increase Jc in Mrt48 and Mht1 wires.
4. Conclusions Basic properties of Ti sheathed MgB2 wires made of oxidized in situ and mechanically alloyed powders have been studied. It was found that oxidation of Mg + B mixture in air is not as risky as for ex situ made MgB2 and decreasing the critical current density only slightly. Surprisingly, the oxidation of intensively milled (mechanically alloyed) powder is improving critical current density especially in high magnetic field (Jc at 10T increased by 47–100%). It enables a safety handling with fine grain mechanically alloyed powder in air. It is also shown that widely used Rowell analysis (for grain connectivity) cannot be used generally for prediction of Jc(B) performance of any kind of MgB2 samples, which are influenced also by the pinning at grain boundaries and at nanosize and well distributed defects.
This work was supported by the Slovak Scientific Agency under the Project APVV-0495-10. References [1] M. Kurth, et al.