High magnetic field enhancement of the critical current density by Ge, GeO2 and Ge2C6H10O7 additions to MgB2

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ScienceDirect Scripta Materialia 82 (2014) 61–64 www.elsevier.com/locate/scriptamat

High magnetic field enhancement of the critical current density by Ge, GeO2 and Ge2C6H10O7 additions to MgB2 D. Batalu,a G. Aldica,b S. Popa,b L. Miu,b M. Enculescu,b R.F. Negrea,b I. Pasukb and P. Badicab,⇑ a

University Politehnica of Bucharest, Splaiul Independentei 313, 060042 Bucharest, Romania b National Institute of Materials Physics, Atomistilor 105 bis, 077125 Magurele, Romania Received 11 February 2014; revised 19 March 2014; accepted 25 March 2014 Available online 1 April 2014

Ge, GeO2 and Ge2C6H10O7 additions to MgB2 obtained by ex situ spark plasma sintering significantly enhance the critical current density Jc in high magnetic fields. A Jc(T = 20 K) of 102 A cm2 is obtained at 3.9 T in the pristine sample and at 5.8 T in the MgB2(Ge2C6H10O7)0.0014 sample. The decrease in the critical temperature for added samples is less than 1 K and Jc(20 K, H = 0) shows a small decrease from 5.5  105 A cm2 in the pristine sample to 3.9  105 A cm2 in MgB2(Ge2C6H10O7)0.0014 sample. Ge does not substitute into the MgB2 lattice. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: MgB2; Ge-based additions; Spark plasma sintering; Critical current density

MgB2 is a promising practical superconductor [1], due to its lightness, low price, wide availability and interesting physics. Most of the metalloid elements or their compounds have been added to MgB2. The most intensively studied are Si and especially SiC. Si [2] was reported to enhance pinning, the irreversibility field (Hirr) and the critical current density (Jc) at low temperatures, without compromising the critical temperature (Tc). SiC was found to be an important addition, promoting a strong enhancement of Jc and Hirr, mostly owing to the effect of substituting C for B [3,4]. According to the Mg–B phase diagram [5], the addition of B to MgB2 will lead to an increase in the amount of MgB4 if the processing parameters are favorable for the appropriate chemical reactions. In this case, MgB4 will degrade the superconducting properties [6]. On the other hand, the formation of higher borides as nanophases have been mentioned as good pinning centers [7,8]. The addition of Sb degrades both the Jc and Hirr [9], while Sb2O3 enhances the Jc at low temperatures and in high magnetic fields, as well as Hirr [9–11]. Te and Te2O3 have been reported to improve the superconducting properties of

⇑ Corresponding

author. Tel.: +40 21 3690170; fax: +40 21 3690177; e-mail: [email protected]

MgB2 for ex situ routes [12]. On the other hand, a neutral influence of Te for in situ routes was shown in Ref. [13]. The best results for ex situ routes were reported for SiC and Te co-doping [14]. However, no data is available for poisonous As-based additives. Vortex pinning interactions in MgB2 take place through dl and dTc mechanisms, while magnetic interactions are negligible [15]. It is considered that the substitution of C for B in the crystal lattice of MgB2 produces disorder and shortens the mean free path l, enhancing the dl pinning. The dTc mechanism is active due to grain boundaries, nanoscale impurities and defect pinning. In this work, the proposed metalloid additions to MgB2 are Ge, GeO2 and Ge2C6H10O7. The first two additions are expected to influence mainly the dTc mechanism, while Ge2C6H10O7 would influence both mechanisms. The result shown in this work is a pinning enhancement, leading to higher Jc and Hirr in the added samples compared to pristine MgB2. It is noteworthy that MgB2 is sensitive not only to additions, but also to processing routes. We applied spark plasma sintering (SPS) to MgB2-based powder mixtures (ex situ route). This technology promotes the formation of dense samples [16], and shows unconventional activation-specific processes [17], owing to the use of a pulsed current. These processes

http://dx.doi.org/10.1016/j.scriptamat.2014.03.024 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

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D. Batalu et al. / Scripta Materialia 82 (2014) 61–64

strongly influence the grain boundaries, which are excellent pinning centers [18]. SPS is also a flexible technique, the high heating and cooling rates avoiding grain coarsening. The achievement of high densities and the possibility of grain boundary and particle-size refinement recommend SPS as a useful method for the preparation of top-quality MgB2. The raw materials (Alfa Aesar) were MgB2 (99.5%, 1–2 lm, decomposition temperature Tdec = 890 °C), Ge (99.999%, 152 lm, melting point Tm = 937 °C), GeO2 (99.999%, 14.3 lm, Tm = 1086 °C), and Ge2C6H10O7 (99.7% purity, Tdec = 320 °C). Powder mixtures of MgB2 and additions with compositions MgB2(Ge)0.005, MgB2(GeO2)0.005 and MgB2(Ge2C6H10O7)0.0025 and 0.0014 were loaded into graphite dies and SPSed (FCT Systeme GmbH – HP D 5, Germany) into samples of 2 cm diameter at 1150 °C for 3 min and under a uniaxial maximum pressure of 95 MPa. The apparent densities of the SPSed samples were measured by Archimedes’ method, in toluene, with a KERN ALT 220-4M density balance (Table 1). The relative density (Table 1) was calculated [19] by considering that the samples contain MgB2 (2.63 g cm3), MgO (3.58 g cm3), MgB4 (2.49 g cm3) and Mg2Ge (3.08 g cm3). The weight fraction of each phase was estimated by Rietveld refinement (MAUD v.2.31 software [20]) of the X-ray diffraction (XRD; Bruker-AXS D8 ˚ ) data. MgB2 crysADVANCE, radiation, k = 1.5406 A tallite size and residual strain were calculated based on the XRD data using the Williamson–Hall procedure [21]. The microstructure was visualized using scanning electron microscopy (SEM; Zeiss EVO50) and transmission electron microscopy (TEM; JEM ARM 200F/ Gatan Quantum SE filter). The temperature variation of the magnetic moment, m, and the relaxation curves, m(t), were measured with an MPMS-7T (Quantum Design). The relaxation time, t, was taken to be zero when the magnet charging was finished. For magnetic hysteresis measurements we used a VSM-9T (Cryogenic) magnetometer. Jc was calculated with the Bean relation [22,23], Jc (A cm2) = 60|m|/V‘ (1), where V is the sample volume (cm3), ‘ is the basal

square side (cm) and, for T well below Tc, m is identified with the irreversible magnetic moment (in emu). The relative density (Table 1) shows high values, between 93.8% and 96.8%, which are not different enough to influence Jc significantly. For pristine MgB2, Jc was found to be constant when the relative density was above 90% [16,24]. Considering the XRD patterns (Supplementary Fig. 1) and the decomposition reaction of MgB2 [2MgB2 ! MgB4 + Mg (2)], the general reactions during SPS of our added samples are: 8MgB2 þ Ge þ O2ðresidual

in SPS chamberÞ

! 4MgB4 þ 2MgO þ Mg2 Ge

ð3Þ

8MgB2 þ GeO2 ! 4MgB4 þ 2MgO þ Mg2 Ge

ð4Þ

8MgB2 þ 0:5Ge2 C6 H10 O7 ! 4MgB4 þ 2MgO þ Mg2 Ge þ 3C þ 2:5H2 þ 0:75O2

ð5Þ

The lattice parameters a and c for the pristine sample “a” and for the samples with Ge (“b”) and GeO2 (“c”) do not change significantly (Table 1). Moreover, in the Ge (“b”) and GeO2 (“c”) added samples the decrease in the midpoint critical temperature (Tc,midpoint) is small when compared to that of the pristine sample “a” (<0.25 K, Supplementary Fig. 2), while the onset critical temperature (Tc,onset) is practically the same (38.8 K). These results suggest that Ge does not substitute B in the MgB2 crystal structure. For the samples “d” and “e”, with C-containing Ge2C6H10O7 added (Table 1), a is smaller, while c is approximately the same as for the pristine sample “a”. This is the typical behavior of C substitution for B [25,26], and is accompanied by the suppression of the transition temperature in pristine MgB2. The highest suppression of Tc,midpoint, of 1 K, is for sample “d”. This sample has the same amount of Ge as the samples with added Ge or GeO2. Sample “e” (containing a smaller amount of C) has a lower suppression (0.8 K), as expected. Thus, it is confirmed that Ge does not substitute B. Based on the relationship between a and the amount of C substituting for B [25,26], in sample “d”, with Ge2C6H10O7 added ˚ ), one has C  0.7 at.%. Some C contami(a = 3.081 A

Table 1. Sample (S) starting composition, apparent and relative density, MgB2 lattice parameters, secondary phases, average calculated crystallite size (D*), and the corresponding residual strain. S

*

Starting composition

“a”

MgB2

“b”

MgB2(Ge)0.005

“c”

MgB2(GeO2)0.005

“d”

MgB2(Ge2C6H10O7)0.0025

“e”

MgB2(Ge2C6H10O7)0.0014

Apparent density (g/cm3)/relative density (%)

Lattice parameter ˚ ) c, (A ˚) a, (A

2.52 95.3 2.59 96.8 2.52 94.4 2.46 93.8 2.50 94.1

3.083 3.524 3.084 3.528 3.084 3.528 3.081 3.527 3.082 3.527

Phase (wt.%) MgB2

MgB4

Average crystallite size, D (nm)/ Residual strain MgO

Mg2Ge

*

MgO

Mg2Ge

250 ± 150 /0.71 252 ± 202 /0.74 223 ± 147 /0.72 183 ± 115 /0.81 192 ± 130 /0.77

159 ± 34

–

75 ± 41

64 ± 25

67 ± 46

61 ± 32

90 ± 49

–

58 ± 40

–

84.8

7.9

7.3

–

79.4

12.9

6.8

0.9

77.3

13.7

8.0

1.0

79.7

9.8

9.5

1.0

81.2

9.9

8.3

0.64

MgB2

Crystallite size of MgB2 is shown for a comparative qualitative analysis; it is well known that Williamson-Hall method is not reliable when crystallite size is above 100–150 nm.

D. Batalu et al. / Scripta Materialia 82 (2014) 61–64 1300

1.00

m(t) / m(t1)

Uc 1100

0.99 0.98 0.97 0.96 0.1

U* (K)

nation (0.42 at.%) occurs in the pristine, Ge- and GeO2added samples due to the use of the graphite die in the SPS processing (a is smaller and c is similar to the values ˚ , c = 3.525 A ˚ [11]). of the raw MgB2 powder, a = 3.088 A The above results suggest that the samples with added Ge (“b”) and GeO2 (“c”) can be considered as typical composites, and the enhancement of Jc in intermediate and high fields or the reduced Jc/Jc0 (Fig. 1) for these samples vs. the pristine sample is mainly the result of the dTc- pinning. The two samples have a similar amount of Mg2Ge (Table 1). A greater amount of MgO and MgB4 and a smaller amount of MgB2 (which are in good agreement with reactions (2) and (4)) were found in sample “c”, with GeO2 added. At the same time, the Jc and Jc/Jc0 vs. H curves at different temperatures (Fig. 1, samples “b” and “c”) are almost identical. This suggests that Ge (through Mg2Ge) has a major contribution to the pinning enhancement. The average crystallite size (Table 1) of Mg2Ge is about 60 nm. The average crystallite size of MgO is in the same range or larger (Table 1). These values are about three times larger than the coherence length n of MgB2. Thus, relatively large impurities cannot be considered as pinning centers, a large contribution to pinning in our samples being supplied by the interfaces between MgB2 and impurities. An unambiguous measure of the characteristic vortex pinning energy is the maximum value Uc [27] observed in the T variation of the normalized vortex-creep activation U* = TDln(t)/Dln(|m|) (Fig. 2). To determine U*, we used the m(t) curves for H = 3 T at various temperatures for t > 100 s (to avoid some of the influence of the microflux jumps at early relaxation stages). The inset of Figure 2 is a comparative example of the relaxation curves (normalized to the first measured m), where a smaller slope in the log–log plot means a higher U*. In Figure 2, Uc is higher for the samples with added Ge (“b”) and GeO2 (“c”) than for the pristine sample “a”, meaning that there is stronger pinning in the added samples. Also, there is a slightly higher Uc for the Ge than for the GeO2 addition. It is possible that this is related to the distribution of the impurity phases (Fig. 3). For the Ge addition, the phases Mg2Ge (Fig. 3(2)) and MgB2 (Fig. 3(1)) are in close proximity, while in the case

63

T = 10 K 0.2

0.5

1

2

3

t (10 s)

900

700

'a' 'b' 'c' 'e' 8

10

30 kOe 12

14

16

18

T (K)

Figure 2. Vortex-creep activation energy U* vs. temperature exhibiting the (maximum) characteristic pinning energy Uc for the SPSed pristine (“a”) and added samples with Ge (“b”), GeO2 (“c”) and Ge2C6H10O7 (“e”). The inset presents an example of two relaxation curves at 10 K for samples added with GeO2 (“c”) and Ge2C6H10O7 (“e”). Zero-fieldcooling conditions were applied with H perpendicular to the largest side of the parallelepipeds (1.5  1.5  1 mm3) cut from the center of the SPSed discs.

Figure 3. SEM and energy-dispersive X-ray spectroscopy maps (for each element and for Mg-B-O-C-Ge) for samples added with Ge (“b”) and GeO2 (“c”). TEM, typical selected area electron diffraction and electron energy loss spectroscopy maps of BK (188 eV), CK (284 eV), OK (532 eV), GeL2,3 (1217 eV) and MgK (1305 eV) edges, and of Mg-BO-C-Ge taken on sample “e” (with added Ge2C6H10O7). Grains or regions ascribed to (1) MgB2, (2) Mg2Ge, (3) MgO and (4) B(C)–O phases are indicated.

Figure 1. Critical current density vs. magnetic field at 20 K and 5 K for samples “a”–“e” (Table 1). For a better comparison, Jc normalized to the self-field critical current density Jc0 is also plotted.

of GeO2, owing to the included oxygen, Mg2Ge is surrounded and isolated by oxygen-rich phases such as MgO (Fig. 3(3)). Thus, fewer superconductor–normal interfaces are available for pinning. As well as interfaces, nanoimpurities of size close to n are also efficient pinning centers. Filamentary grains with Ge and Mg (ascribed to Mg2Ge) of 10–25 nm length and 2–3 nm

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D. Batalu et al. / Scripta Materialia 82 (2014) 61–64

thickness are visualized in Figure 3. For the ex situ method, Birajdar et al. [28] noted that boron-rich phases are not easily detected by XRD, though they do occur and can improve pinning. Boron-rich phases (such as the higher borides of magnesium, e.g. MgB7 or MgB20 [29,30]) are stable in the Mg–B system under our SPS conditions. The occurrence of these phases may explain the slightly smaller amount of MgB4 than expected from reactions (2)–(5) in the Ge2C6H10O7-added samples “d” and “e”, when compared with the Ge- and GeO2-added samples “b” and “c” (Table 1). However, another explanation, such as different Mg evaporation during SPS for samples with different additives, is possible. Grains ascribed to B(C)–O phases (likely B2O3 in Fig. 3(4)) were also observed. Unreacted GeO2 was detected in sample “c”, with added GeO2. All these residual phases can influence pinning, as well. The samples with added Ge2C6H10O7 take advantage not only of a dTc but also of the dl pinning mechanism, owing to the substitution C for B. The effect of C substitution is accompanied by a substantial increase in residual strain (Table 1). This residual strain increase follows the pinning enhancement [31], i.e. the increase in Uc (Fig. 2). However, the residual strain does not follow the increase in Jc for samples “d” and “e” (Fig. 1), with added Ge2C6H10O7. The result points to opposite effects influencing Jc for complex additions, such as Ge2C6H10O7, and this is of interest for further research. It is remarkable that sample “e”, which contains less addition, i.e. less Ge than the optimally Ge-doped sample “b”, less C than sample “d” and the same amount of oxygen as the optimally GeO2-doped sample “c”, has the biggest Jc and Jc/Jc0 in intermediate and high fields (Fig. 1). The maximum Jc in a high field and at 20 K is competitive. For ex situ processing, a Jc of 102 A cm2, at 20 K, is attained at 5.8 T in sample “e” from this work and at 5.3 T for the co-addition of SiC and Te [14]. For in situ processing, some of the best reported values are 6.3 T for MgB2 added with 10% Ti [7], 6.5 T for MgB2 added with 2% graphene oxide [32] and 10 T for MgB2 added with 3.5% C [7], though the latter was obtained under special high-pressure processing conditions (2 GPa). It is also interesting that the self-field Jc0 showed only a relatively small decrease for our addition samples. At 20 K, the decrease is from Jc0 = 5.5  105 A cm2 for the pristine sample “a” to Jc0 = 3.9  105 A cm2 for sample “e”, with added Ge2C6H10O7. For example, in Ref. [32], which shows simultaneous enhancement of the Jc in both a zero field and a high field in graphene oxide-added MgB2 in situ samples, Jc0(20 K) is 6.35  105 A cm2. D.B. and R.F.N. acknowledge POSDRU/89/ 1.5/S/54785 and PN-II-ID-PCE-2012-1-0362, respectively. Work at NIMP was supported by Core Program PN09-450, Romania.

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