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In situ and ex situ Cu doping of MgB2
Physica C 418 (2005) 99–106 www.elsevier.com/locate/physc
In situ and ex situ Cu doping of MgB2 S.K. Chen b
a,b,* ,
M. Majoros b, J.L. MacManus-Driscoll
a,b
, B.A. Glowacki
a,b
a Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, UK Interdisciplinary Research Centre in Superconductivity, University of Cambridge, Madingley Road, Cambridge CB3 0HE, UK
Received 13 September 2004; accepted 14 November 2004 Available online 25 December 2004
Abstract The effect of in situ and ex situ Cu doping on MgB2 was studied in the annealing temperature range of 600 °C– 900 °C. Attempts to substitute Cu into MgB2 resulted in the formation of Mg–Cu phases. Within the limits of error, neither lattice strain nor significant alteration in a- and c-axes was introduced by varying annealing temperature. This indicates that all the samples showed negligible Cu substitution or Mg deficiency. Both in situ and ex situ Cu doping caused an increase of density in the samples annealed at 800 °C by 28% and 26%, respectively. SEM micrographs show that the in situ samples consist of larger grains with low Cu content as well as agglomerates of fine grains which correspond to Cu rich regions. On the other hand, compositional variation in the ex situ sample is less pronounced. The reaction temperature causes only minor changes in the Tc onset but influences the superconducting volume fraction. Ó 2004 Elsevier B.V. All rights reserved. PACS: 74.62.Bf; 74.62.Dh; 74.70.Ad Keywords: MgB2; In situ and ex situ Cu doping; Annealing temperature
1. Introduction In response to the unexpected discovery of Tc 40 K in MgB2, a great deal of work has been carried out of doping on Mg or B. Several theoretical calculations predict that electron doping is unfavourable while hole doping or isoelectronic substitution may improve Tc [1]. By electron dop*
Corresponding author. Tel.: +44 1223 767140; fax: +44 1223 337074. E-mail address: [email protected] (S.K. Chen).
ing, Tc would be degraded due to the decrease in the density of states. However, experimental results have shown that both hole and electron doping are detrimental to Tc [2–5]. In most cases, Tc drops when doping produces substitution or remains unchanged when substitution does not take place [6]. In this regard, the drop in Tc is attributed to the change in lattice constant or strain [2–5,7]. So far, there is no consensus on whether the aor c-axis lengths can be directly linked to Tc depression. By Mn and Al doping, a contraction in c-axis is observed but there is no significant
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S.K. Chen et al. / Physica C 418 (2005) 99–106
alteration to the a-axis [2,3,8]. Conversely, Li and C doping leads to a-axis contraction but leaves caxis unchanged [4,5]. Both contraction in the aand c-axes can lead to degradation of Tc. In fact, it has been established in the high pressure studies that compression of lattice constants has a negative effect on Tc pointing towards BCS-type superconductivity in MgB2 [9]. This is in contrast to hole superconductivity proposed by Hirsch (pressure reduces the intraplane B–B distance) [10]. Degradation of Tc can also be seen in structurally disordered MgB2 caused by ion irradiation [11]. In Mg-deficient Mg1xB2 the a-axis decreases while the c-axis increases linearly with x but Tc did not change remarkably [12]. Recently, some interest has been focussed on Cu doped MgB2 [13,14]. Kazakov et al. reported that Cu can hardly enters the lattice structure of MgB2 [13]. They did not observe any change in lattice parameters or Tc onset though there was a broadening in transition temperature resulting from the formation of Mg–Cu phases. By mixing CuB24 and Mg, Tampieri et al. found a maximum of 3 mol% Cu substituted for Mg [14]. Their samples also contained multiple phases of Cu2y Mg1+y and MgB4. Lattice parameters in both the a- and c-axes decreased with increasing Cu content while Tc decreased by only 0.5 K. It is believed the decrease in Tc could result from lattice contraction which gives rise to hole transfer from the planar px,y band to the three-dimensional pz band. This increases the number of holes in the boron plane and hence decreases Tc. In the early stages of the discovery of MgB2, there was a report of a 49 K Tc in Cu-doped MgB2 but the claim was later withdrawn [15]. Apart from this, Majoros et al. observed an anomalous AC susceptibility decrease at around 50 K in Cu sheathed powder-in-tube MgB2 wires [16]. The authors related this phenomenon to the variation of Cu substitution in MgB2. On the other hand, density functional calculations show that Tc as high as 50 K is achievable with full replacement of Mg by Cu obtaining CuB2 [17]. However, a preferred phase of MgCu2 rather than CuB2 will be produced by substituting Cu into MgB2 and CuB2 is not known to be an equilibrium phase. The difficulty in synthesising Cu substituted
MgB2 is due to the high affinity between Cu and Mg that leads to the formation of Mg–Cu compounds, especially MgCu2. In this work, we study the substitution of Mg in MgB2 via in situ and ex situ Cu doping. The influence of Cu doping on MgB2 at different annealing temperatures is discussed based on results from phase formation and superconducting properties as well from microanalysis.
2. Experimental The starting powders for sample preparation were crystallines MgB2 (98%, 325 mesh), Mg (99.8%, 325 mesh), B (98%, 325 mesh) and Cu (99.5%, 150 mesh) from Alfa Aesar. Two set of samples were prepared using the conventional solid state reaction according to the stoichiometries: 1. 0.8Mg + 0.2Cu + 2B (in situ) 2. MgB2 + 0.2Cu (ex situ) First, appropriate amounts of MgB2, Mg, B and Cu were weighed, then well mixed and ground for 1 h in an agate mortar. This was followed by pressing the resultant powder into pellets of 13 mm in diameter and about 2–3 mm thick using a hydraulic presser with the applied load of 6 tonnes. The pellets were then encapsulated in separate quartz tubes, evacuated to 106 torr and put into a furnace for annealing at heating and cooling rates of 20 °C/min. Annealing was conducted in the temperature range of 600 °C–900 °C and maintained at the temperature for 1 h. The furnace was then switched off and the samples allowed to cool down to room temperature. X-ray powder diffraction in the step-scanning mode h–2h with 0.05° increments was performed in a Philips PW1050 diffractometer with CuKa radiation ˚ ). The onset of superconductsource (k = 1.5418 A ing temperature, Tc was determined using LakeShore 7000 AC susceptometer. The samples were first zero field cooled and measured in the field warming condition. The microstructures and composition analysis of the samples were carried out using JEOL 5800 LV scanning electron micro-
S.K. Chen et al. / Physica C 418 (2005) 99–106
scope equipped with Energy-Dispersive Spectrometer (EDS).
3. Results and discussion 3.1. Phase formation Fig. 1(a) and (b) show the X-ray diffraction patterns of the in situ and ex situ Cu doped samples, respectively. The competition in reaction among Mg, Cu and B and between MgB2 and Cu has led to different phase formations in the in situ and ex situ samples. It is known from the phase diagrams that Cu has a limited solubility of only 0.013 at% in Mg and the solubility limit of B in
(a)
101
Cu is about 0.3 at% [18]. More impurities are formed at higher annealing temperature as shown in Table 1. MgO was found to be present in the in situ samples and there was already some MgO in the commercial MgB2 powder. CuB24 is the only compound produced by reaction of Cu and B as its existence can be seen in all the in situ but only in some of the ex situ samples. Multiple phases were observed in the in situ samples even at low annealing temperatures. MgB2 was formed at 600 °C, which is still below the melting point of magnesium. In addition, the congruently melted compounds of Mg2Cu and MgCu2 are also formed. Unlike the MgB2 compound, the elemental form of Mg is more reactive and hence gives rise to vigorous reaction with Cu. Mg2Cu is not stable at higher annealing temperatures as this phase was not present above 600 °C. Thermodynamically, the Gibbs free energy of MgCu2 is lower than that of Mg2Cu and this explains why the formation of MgCu2 is more favourable [19]. The volatility of Mg at 900 °C resulted in Mg deficiency and favoured the formation of secondary phases such as MgB4 and MgB6. The formation of MgB6 at 900 °C in the in situ sample may imply that Mg deficiency is more severe than in the ex situ sample. It is noticed that the formation of MgCu2 is always accompanied by CuB24 in both in situ and ex situ samples. For ex situ samples, no reaction between Cu and MgB2 was observed in the temperature range Table 1 Phase formation in the in situ and ex situ samples at different annealing temperatures
(b) Fig. 1. X-ray powder diffraction patterns of (a) in situ and (b) ex situ Cu doped samples annealed at 600 °C–900 °C.
Annealing temperature
In situ
Ex situ
600 °C
MgB2, Mg2Cu, CuB24 MgO, MgCu2
MgB2, Cu, MgO
700 °C
MgB2, MgCu2, CuB24 MgO
MgB2, Cu, MgO
800 °C
MgB2, MgCu2, CuB24 MgO
MgB2, MgCu2, CuB24 MgB4 MgO
900 °C
MgB2, MgCu2, CuB24 MgB4 MgB6 MgO
MgB2, MgCu2, CuB24 MgB4 MgO
S.K. Chen et al. / Physica C 418 (2005) 99–106
3.2. Microstructures and compositional analysis The density of the Cu doped samples is shown in Fig. 4. As also included in Fig. 4, the density of the pure in situ and ex situ MgB2 annealed at
3.120 Cu doped (in-situ)
3.110
Cu doped (ex-situ) pure sample (in-situ)
3.100 a -axis (Å)
of 600 °C–700 °C which is below the melting temperature of Cu (1085 °C) and decomposition temperature of MgB2 (800 °C). At 800 °C–900 °C, some MgB2 started to decompose into MgB4 but no MgB6 was detected from X-ray. Compared to in situ samples, MgB4 is only formed at 900 °C. The formation of MgCu2 also indicates that higher annealing temperatures are necessary for ex situ reaction between Cu and MgB2. The cell parameters obtained from Rietveld refinements (Philips XÕ Plus) are displayed in Fig. 2. The relatively large error in the in situ sample annealed at 900 °C is due to its poor quality. The MgB2 peaks in this sample are much weaker than those annealed at lower temperatures. Besides, some MgB2 peaks also overlap with impurity phases. The lattice parameters of both in situ and ex situ pure MgB2 prepared at 800 °C are also included in Fig. 2 for comparison. Within the error of calculation, the a- and c-axes did not change with increasing annealing temperature. Fig. 3 shows the strain development by Cu doping estimated by using a WilliamsonHall plot [20]. We could not obtain the lattice strain from the in situ sample annealed at 900 °C. The weak diffraction peaks caused nonlinearity in the Williamson-Hall plot making estimation of strain difficult. From Fig. 3, we see no systematic change of lattice strain with annealing temperature in the ex situ sample though there is a slight increase in the in situ sample annealed above 600 °C. Since there is no appreciable change in both lattice constant and strain, our samples are not successfully doped with Cu. This also supports previous finding that Cu hardly substitutes for Mg in MgB2 [13,14]. The annealing temperature only caused different phase formation and did not result in Mg deficiency to any great extent. As reported by Xiao et al. [12], the reduction in nominal Mg content does not cause a serious Mg deficiency.
pure sample (ex-situ)
3.090 3.080 3.070 3.060 3.050 3.040 550
(a)
650
750
850
950
Annealing temperature (˚C) 3.600 Cu doped (in-situ)
3.580
Cu doped (ex-situ)
3.560
pure sample (ex-situ)
pure sample (in-situ)
c -axis (Å)
102
3.540 3.520 3.500 3.480 3.460 550
(b)
650
750
850
950
Annealing temperature (˚C)
Fig. 2. Lattice parameters of (a) a-axis and (b) c-axis for in situ and ex situ samples annealed at 600 °C–900 °C.
800 °C are 1.05 g/cm3 and 1.37 g/cm3, respectively. Cu doping enhanced the density of both in situ and ex situ samples annealed at 800 °C by 28% and 26%. In situ Cu doping in the annealing temperature range of 600 °C–900 °C led to a sample density of 1.49 g/cm3–1.79 g/cm3. For the ex situ sample, the density increased from 1.64 g/cm3 to 1.77 g/cm3 in the same temperature range. Relatively, the increment of sample density is higher in the in situ samples. However, these values are still low compared to the theoretical value (2.6 g/ cm3).
S.K. Chen et al. / Physica C 418 (2005) 99–106
103
0.30 in situ 0.25
Relative strain (%)
ex situ 0.20
0.15
0.10
0.05
0.00 550
650
750
850
950
Annealing temperature (˚C)
Fig. 3. Development of strain with annealing temperature.
2.50
Density (g/cm3)
2.00
1.50
1.00 doped sample (in situ)
0.50
doped sample (ex situ) pure sample (in situ) pure sample (ex situ)
0.00 550
650
750
850
950
Annealing temperature (˚C)
Fig. 4. Annealing temperature dependence of sample density.
It can be seen from the SEM micrographs in Fig. 5 that both in situ and ex situ samples annealed at 800 °C are porous. The in situ sample consists of larger grains (regions 3 and 5) with particle size ranges from a few lm to as large as about 15 lm and agglomerates of fine grains (are apparent regions 1, 2, 4 and 6). As shown in Table 2, the larger grains contain very low Cu but large amounts of B. Therefore, these regions can be considered as MgB2 and CuB24 rich. For regions with
Fig. 5. Scanning electron micrographs of (a) in situ and (b) ex situ Cu doped samples annealed at 800 °C.
fine grains, the Cu content is substantially higher. Although the amount of B is higher than Mg, this is not as great as found in larger grains. In regions where Mg has a higher proportion than B (region 2), the content of Cu is lower than in those where the amount of Mg is lower than B (regions 1 and 4). In region 6, there is only 3 at% of Cu, however, a significant amount of oxygen was also found there. This shows that the fine grains are either Cu or oxygen rich, suggestive of MgCu2 and MgO.
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Table 2 Results from EDS for the in situ and ex situ Cu doped samples Regions
1
2
3
4
5
6
In situ Mg (at%) B (at%) O (at%) Cu (at%)
6.68 30.68 1.9 60.74
36.47 26.14 2.65 34.74
10.67 74.49 5.06 9.77
3.35 29.87 3.77 63.01
14.52 84.04 0.94 0.5
18.15 61.3 17.54 3
30.19 55.08 12.06 2.67
27.22 58.48 9.46 4.84
31.8 53.94 6.5 7.76
45.45 49.92 3.97 0.66
41.9 38.56 6.82 12.72
27.64 51.38 10.57 10.41
Total (100 at%) Ex situ Mg (at%) B (at%) O (at%) Cu (at%) Total (100 at%)
SEM images in Fig. 5b show that the ex situ sample have randomly distributed fine grains with sizes as large as a few lm and some of them are as small as hundreds of nm. From the EDS results, the Mg content in regions 1–6 is either lower than or comparable to B. This indicates that reaction between MgB2 and Cu is not vigorous, as much of the MgB2 still remained. The oxygen and Cu content in these samples shows no significant difference except in region 4 where Cu content is low. Relatively, the ex situ samples have lower Cu but higher oxygen content than the in situ samples. The higher amount of oxygen is probably due to MgO contamination in the commercial MgB2 powder. 3.3. Superconducting properties Fig. 6 show the results from AC susceptibility measurements in the temperature range of 5– 65 K. The Tc onset which was determined from the deviation point of transition curve is shown in Fig. 7. Tc is higher in the in situ samples than in the ex situ samples. The Tc of both in situ and ex situ samples only shows minor changes within the annealing temperature range of 600 °C– 900 °C. The variation of superconducting transition temperature, DT for in situ and ex situ samples is 0.5 K and 1.2 K, respectively. For in situ samples, the transition is sharp except for the sample annealed at 900 °C in which Tc could not be detected as the diamagnetic signal
is weak (Fig. 6a). There is no significant change of superconducting volume fraction in the annealing temperature range of 600 °C–800 °C. The small variation in vint can be accounted for by calculation error. The large magnitude of susceptibility (<8) indicates the bulk superconductivity in these samples. At 900 °C, the drastic drop of the superconducting volume fraction mainly resulted from a decrease in the phase volume of MgB2 by forming more secondary phases such as MgB4, MgB6 and MgCu2. Tc only showed a slight increase from 38.0 K at 600 °C up to 38.5 °C at 700 °C before it decreased slightly to 38.4 K at 800 °C. As shown in Fig. 6(b), broadening in the transition temperature is observed in the ex situ samples. Besides, the samples annealed at 800 °C–900 °C exhibit a dual-transition indicating the presence of multiple superconducting phases. Tc was increased from 37.5 K at 600 °C up to 38.3 K at 800 °C before it dropped to 37.1 K at 900 °C. This is accompanied by the same evolution of superconducting volume fraction which increased from 600 °C to 800 °C and decreased at 900 °C (jvint800 °Cj > jvint700 °Cj > j vint600 °Cj > jvint900 °Cj).
4. Summary In situ and ex situ Cu doping in MgB2 has been carried out in the temperature range of 600 °C–
S.K. Chen et al. / Physica C 418 (2005) 99–106 0.0
105
38.6 -0.0010
Susceptibility,χ int
-0.1
-0.3
38.2
-0.0025 -0.0030
38.0
-0.0035
600 ˚C
-0.4
700 ˚C -0.5
-0.0040
T c (K)
Susceptibility,χ int
-0.2
38.4
-0.0015 -0.0020
-0.0045 5
15 25 35 Temperature (K)
45
800 ˚C
-0.6
900 ˚C
37.8 37.6 37.4
-0.7
37.2 in situ
-0.8
37.0
ex-situ
-0.9 5
15
(a)
25
35
45
0.0
700 ˚C
900 ˚C Susceptibility,χ int
750 850 Annealing temperature (˚C)
950
800 ˚C
-0.2 -0.3 -0.4 -0.5 -0.6 -0.7 -0.8 5
(b)
650
Fig. 7. Variation of Tc onset with annealing temperature in in situ and ex situ samples. Solid lines are the guides for the eyes.
600 ˚C
-0.1
36.8 550
Temperature (K)
15
25
35
45
For both in situ and ex situ samples, Cu doping leads to an increase of density in sample annealed at 800 °C by 28% and 26%, respectively. SEM micrographs show that the in situ samples consist of larger grains with low Cu content and agglomerates of fine grains corresponds to Cu rich regions. On the other hand, compositional variation in the ex situ sample is less pronounced. The increasing annealing temperature causes only minor changes in the Tc onset but decreases the superconducting volume fraction.
Temperature (K)
Fig. 6. AC susceptibility measured at applied AC field of 1 Oerms and 333.3 Hz for (a) in situ and (b) ex situ Cu doped samples. The inset shows the sample annealed at 900 °C.
900 °C. Attempts to substitute Cu into MgB2 resulted in the formation Mg–Cu phases but showed no evidence of Cu doping into the lattice structure. Within the limits of error, no significant lattice strain nor alteration in a- and c-axes was introduced by varying annealing temperature. This indicates that all the samples suffered only minor Mg deficiency. At higher annealing temperatures (P800 °C), the formation of MgB4 and MgB6 is probably due to a loss of Mg to Mg–Cu alloys.
Acknowledgments We would like to thank R. Stern and D. Smith for the technical assistance. S.K. Chen acknowledges UPM and EPSRC for the financial support.
References [1] N.I. Medvedeva, A.L. Ivanovskii, J.E. Medvedeva, A.J. Freeman, Phys. Rev. B 64 (2001) 020502R. [2] S. Xu, Y. Moritomo, K. Kato, J. Phys. Soc. Jpn. 70 (2001) 1889. [3] J.S. Slusky, N. Rogado, K.A. Regan, M.A. Hayward, P. Khalifah, T. He, K. Inumaru, S.M. Loureiro, M.K. Haas, H.W. Zandbergen, R.J. Cava, Nature 410 (2001) 343.
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[4] Y.G. Zhao, X.P. Zhang, P.T. Qiao, H.T. Zhang, S.L. Jia, B.S. Cao, M.H. Zhu, Z.H. Han, X.L. Wang, B.L. Gu, Physica C 361 (2001) 91. [5] M. Paranthaman, J.R. Thompson, D.K. Christen, Physica C 355 (2001) 1. [6] I. Felner, Physica C 353 (2001) 11. [7] A. Serquis, Y.T. Zhu, E.J. Peterson, J.Y. Coulter, D.E. Peterson, F.M. Mueller, Appl. Phys. Lett. 79 (2001) 4399. [8] A. Berenov, A. Serquis, X.Z. Liao, Y.T. Zhu, D.E. Peterson, Y. Bugoslavsky, K.A. Yates, M.G. Blamire, L.F. Cohen, J.L. MacManus-Driscoll, Supercond. Sci. Technol. 17 (2004) 1093. [9] J. Tang, L.C. Qin, A. Matsushita, Y. Takano, K. Togano, H. Kito, H. Ihara, Phys. Rev. B. 64 (2001) 132509. [10] J.E. Hirsch, Phys. Lett. A. 282 (2001) 392. [11] Y. Bugoslavsky, L.F. Cohen, G.K. Perkins, M. Polichetti, T.J. Tate, R. Gwilliam, A.D. Caplin, Nature 411 (2001) 561. [12] H. Xiao, W. Peng, W.H. Song, R.C. Ma, L. Zhang, J.J. Du, Y.P. Sun, Physica C 386 (2003) 648.
[13] S.M. Kazakov, M. Angst, J. Karpinski, I.M. Fita, R. Puzniak, Solid State Commun. 119 (2001) 1. [14] A. Tampieri, G. Celotti, S. Sprio, D. Rinaldi, G. Barucca, R. Caciuffo, Solid State Commun. 121 (2002) 497. [15] Y.P. Sun, W.H. Song, J.M. Dai, B. Zhao, J.J. Du, H.H. Wen, Z.H. Zhao, cond-mat/0103101 (2001). [16] M. Majoros, B.A. Glowacki, M.E. Vickers, Supercond. Sci. Technol. 15 (2002) 269. [17] M.J. Mehl, D.A. Papaconstantopoulos, D.J. Singh, Phys. Rev. B 64 (2001) 140509R. [18] T. Massalski, Binary Alloy Phase Diagram, A.S.M. International, Materials Park, OH, 1990. [19] I. Barin, Thermochemical Data of Pure Substances, VCH, Weinheim, 1995. [20] G. Kimmel, D. Dayan, in: R.L. Snyder, J. Fiala, H.J. Bunge (Eds.), Defect and Microstructure Analysis by Diffraction: X-ray Diffraction Broadening Effects in Materials Characterization, Oxford University Press, Oxford, 1999, p. 698.
In situ and ex situ Cu doping of MgB2 S.K. Chen b
a,b,* ,
M. Majoros b, J.L. MacManus-Driscoll
a,b
, B.A. Glowacki
a,b
a Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, UK Interdisciplinary Research Centre in Superconductivity, University of Cambridge, Madingley Road, Cambridge CB3 0HE, UK
Received 13 September 2004; accepted 14 November 2004 Available online 25 December 2004
Abstract The effect of in situ and ex situ Cu doping on MgB2 was studied in the annealing temperature range of 600 °C– 900 °C. Attempts to substitute Cu into MgB2 resulted in the formation of Mg–Cu phases. Within the limits of error, neither lattice strain nor significant alteration in a- and c-axes was introduced by varying annealing temperature. This indicates that all the samples showed negligible Cu substitution or Mg deficiency. Both in situ and ex situ Cu doping caused an increase of density in the samples annealed at 800 °C by 28% and 26%, respectively. SEM micrographs show that the in situ samples consist of larger grains with low Cu content as well as agglomerates of fine grains which correspond to Cu rich regions. On the other hand, compositional variation in the ex situ sample is less pronounced. The reaction temperature causes only minor changes in the Tc onset but influences the superconducting volume fraction. Ó 2004 Elsevier B.V. All rights reserved. PACS: 74.62.Bf; 74.62.Dh; 74.70.Ad Keywords: MgB2; In situ and ex situ Cu doping; Annealing temperature
1. Introduction In response to the unexpected discovery of Tc 40 K in MgB2, a great deal of work has been carried out of doping on Mg or B. Several theoretical calculations predict that electron doping is unfavourable while hole doping or isoelectronic substitution may improve Tc [1]. By electron dop*
Corresponding author. Tel.: +44 1223 767140; fax: +44 1223 337074. E-mail address: [email protected] (S.K. Chen).
ing, Tc would be degraded due to the decrease in the density of states. However, experimental results have shown that both hole and electron doping are detrimental to Tc [2–5]. In most cases, Tc drops when doping produces substitution or remains unchanged when substitution does not take place [6]. In this regard, the drop in Tc is attributed to the change in lattice constant or strain [2–5,7]. So far, there is no consensus on whether the aor c-axis lengths can be directly linked to Tc depression. By Mn and Al doping, a contraction in c-axis is observed but there is no significant
0921-4534/$ - see front matter Ó 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.physc.2004.11.011
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S.K. Chen et al. / Physica C 418 (2005) 99–106
alteration to the a-axis [2,3,8]. Conversely, Li and C doping leads to a-axis contraction but leaves caxis unchanged [4,5]. Both contraction in the aand c-axes can lead to degradation of Tc. In fact, it has been established in the high pressure studies that compression of lattice constants has a negative effect on Tc pointing towards BCS-type superconductivity in MgB2 [9]. This is in contrast to hole superconductivity proposed by Hirsch (pressure reduces the intraplane B–B distance) [10]. Degradation of Tc can also be seen in structurally disordered MgB2 caused by ion irradiation [11]. In Mg-deficient Mg1xB2 the a-axis decreases while the c-axis increases linearly with x but Tc did not change remarkably [12]. Recently, some interest has been focussed on Cu doped MgB2 [13,14]. Kazakov et al. reported that Cu can hardly enters the lattice structure of MgB2 [13]. They did not observe any change in lattice parameters or Tc onset though there was a broadening in transition temperature resulting from the formation of Mg–Cu phases. By mixing CuB24 and Mg, Tampieri et al. found a maximum of 3 mol% Cu substituted for Mg [14]. Their samples also contained multiple phases of Cu2y Mg1+y and MgB4. Lattice parameters in both the a- and c-axes decreased with increasing Cu content while Tc decreased by only 0.5 K. It is believed the decrease in Tc could result from lattice contraction which gives rise to hole transfer from the planar px,y band to the three-dimensional pz band. This increases the number of holes in the boron plane and hence decreases Tc. In the early stages of the discovery of MgB2, there was a report of a 49 K Tc in Cu-doped MgB2 but the claim was later withdrawn [15]. Apart from this, Majoros et al. observed an anomalous AC susceptibility decrease at around 50 K in Cu sheathed powder-in-tube MgB2 wires [16]. The authors related this phenomenon to the variation of Cu substitution in MgB2. On the other hand, density functional calculations show that Tc as high as 50 K is achievable with full replacement of Mg by Cu obtaining CuB2 [17]. However, a preferred phase of MgCu2 rather than CuB2 will be produced by substituting Cu into MgB2 and CuB2 is not known to be an equilibrium phase. The difficulty in synthesising Cu substituted
MgB2 is due to the high affinity between Cu and Mg that leads to the formation of Mg–Cu compounds, especially MgCu2. In this work, we study the substitution of Mg in MgB2 via in situ and ex situ Cu doping. The influence of Cu doping on MgB2 at different annealing temperatures is discussed based on results from phase formation and superconducting properties as well from microanalysis.
2. Experimental The starting powders for sample preparation were crystallines MgB2 (98%, 325 mesh), Mg (99.8%, 325 mesh), B (98%, 325 mesh) and Cu (99.5%, 150 mesh) from Alfa Aesar. Two set of samples were prepared using the conventional solid state reaction according to the stoichiometries: 1. 0.8Mg + 0.2Cu + 2B (in situ) 2. MgB2 + 0.2Cu (ex situ) First, appropriate amounts of MgB2, Mg, B and Cu were weighed, then well mixed and ground for 1 h in an agate mortar. This was followed by pressing the resultant powder into pellets of 13 mm in diameter and about 2–3 mm thick using a hydraulic presser with the applied load of 6 tonnes. The pellets were then encapsulated in separate quartz tubes, evacuated to 106 torr and put into a furnace for annealing at heating and cooling rates of 20 °C/min. Annealing was conducted in the temperature range of 600 °C–900 °C and maintained at the temperature for 1 h. The furnace was then switched off and the samples allowed to cool down to room temperature. X-ray powder diffraction in the step-scanning mode h–2h with 0.05° increments was performed in a Philips PW1050 diffractometer with CuKa radiation ˚ ). The onset of superconductsource (k = 1.5418 A ing temperature, Tc was determined using LakeShore 7000 AC susceptometer. The samples were first zero field cooled and measured in the field warming condition. The microstructures and composition analysis of the samples were carried out using JEOL 5800 LV scanning electron micro-
S.K. Chen et al. / Physica C 418 (2005) 99–106
scope equipped with Energy-Dispersive Spectrometer (EDS).
3. Results and discussion 3.1. Phase formation Fig. 1(a) and (b) show the X-ray diffraction patterns of the in situ and ex situ Cu doped samples, respectively. The competition in reaction among Mg, Cu and B and between MgB2 and Cu has led to different phase formations in the in situ and ex situ samples. It is known from the phase diagrams that Cu has a limited solubility of only 0.013 at% in Mg and the solubility limit of B in
(a)
101
Cu is about 0.3 at% [18]. More impurities are formed at higher annealing temperature as shown in Table 1. MgO was found to be present in the in situ samples and there was already some MgO in the commercial MgB2 powder. CuB24 is the only compound produced by reaction of Cu and B as its existence can be seen in all the in situ but only in some of the ex situ samples. Multiple phases were observed in the in situ samples even at low annealing temperatures. MgB2 was formed at 600 °C, which is still below the melting point of magnesium. In addition, the congruently melted compounds of Mg2Cu and MgCu2 are also formed. Unlike the MgB2 compound, the elemental form of Mg is more reactive and hence gives rise to vigorous reaction with Cu. Mg2Cu is not stable at higher annealing temperatures as this phase was not present above 600 °C. Thermodynamically, the Gibbs free energy of MgCu2 is lower than that of Mg2Cu and this explains why the formation of MgCu2 is more favourable [19]. The volatility of Mg at 900 °C resulted in Mg deficiency and favoured the formation of secondary phases such as MgB4 and MgB6. The formation of MgB6 at 900 °C in the in situ sample may imply that Mg deficiency is more severe than in the ex situ sample. It is noticed that the formation of MgCu2 is always accompanied by CuB24 in both in situ and ex situ samples. For ex situ samples, no reaction between Cu and MgB2 was observed in the temperature range Table 1 Phase formation in the in situ and ex situ samples at different annealing temperatures
(b) Fig. 1. X-ray powder diffraction patterns of (a) in situ and (b) ex situ Cu doped samples annealed at 600 °C–900 °C.
Annealing temperature
In situ
Ex situ
600 °C
MgB2, Mg2Cu, CuB24 MgO, MgCu2
MgB2, Cu, MgO
700 °C
MgB2, MgCu2, CuB24 MgO
MgB2, Cu, MgO
800 °C
MgB2, MgCu2, CuB24 MgO
MgB2, MgCu2, CuB24 MgB4 MgO
900 °C
MgB2, MgCu2, CuB24 MgB4 MgB6 MgO
MgB2, MgCu2, CuB24 MgB4 MgO
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3.2. Microstructures and compositional analysis The density of the Cu doped samples is shown in Fig. 4. As also included in Fig. 4, the density of the pure in situ and ex situ MgB2 annealed at
3.120 Cu doped (in-situ)
3.110
Cu doped (ex-situ) pure sample (in-situ)
3.100 a -axis (Å)
of 600 °C–700 °C which is below the melting temperature of Cu (1085 °C) and decomposition temperature of MgB2 (800 °C). At 800 °C–900 °C, some MgB2 started to decompose into MgB4 but no MgB6 was detected from X-ray. Compared to in situ samples, MgB4 is only formed at 900 °C. The formation of MgCu2 also indicates that higher annealing temperatures are necessary for ex situ reaction between Cu and MgB2. The cell parameters obtained from Rietveld refinements (Philips XÕ Plus) are displayed in Fig. 2. The relatively large error in the in situ sample annealed at 900 °C is due to its poor quality. The MgB2 peaks in this sample are much weaker than those annealed at lower temperatures. Besides, some MgB2 peaks also overlap with impurity phases. The lattice parameters of both in situ and ex situ pure MgB2 prepared at 800 °C are also included in Fig. 2 for comparison. Within the error of calculation, the a- and c-axes did not change with increasing annealing temperature. Fig. 3 shows the strain development by Cu doping estimated by using a WilliamsonHall plot [20]. We could not obtain the lattice strain from the in situ sample annealed at 900 °C. The weak diffraction peaks caused nonlinearity in the Williamson-Hall plot making estimation of strain difficult. From Fig. 3, we see no systematic change of lattice strain with annealing temperature in the ex situ sample though there is a slight increase in the in situ sample annealed above 600 °C. Since there is no appreciable change in both lattice constant and strain, our samples are not successfully doped with Cu. This also supports previous finding that Cu hardly substitutes for Mg in MgB2 [13,14]. The annealing temperature only caused different phase formation and did not result in Mg deficiency to any great extent. As reported by Xiao et al. [12], the reduction in nominal Mg content does not cause a serious Mg deficiency.
pure sample (ex-situ)
3.090 3.080 3.070 3.060 3.050 3.040 550
(a)
650
750
850
950
Annealing temperature (˚C) 3.600 Cu doped (in-situ)
3.580
Cu doped (ex-situ)
3.560
pure sample (ex-situ)
pure sample (in-situ)
c -axis (Å)
102
3.540 3.520 3.500 3.480 3.460 550
(b)
650
750
850
950
Annealing temperature (˚C)
Fig. 2. Lattice parameters of (a) a-axis and (b) c-axis for in situ and ex situ samples annealed at 600 °C–900 °C.
800 °C are 1.05 g/cm3 and 1.37 g/cm3, respectively. Cu doping enhanced the density of both in situ and ex situ samples annealed at 800 °C by 28% and 26%. In situ Cu doping in the annealing temperature range of 600 °C–900 °C led to a sample density of 1.49 g/cm3–1.79 g/cm3. For the ex situ sample, the density increased from 1.64 g/cm3 to 1.77 g/cm3 in the same temperature range. Relatively, the increment of sample density is higher in the in situ samples. However, these values are still low compared to the theoretical value (2.6 g/ cm3).
S.K. Chen et al. / Physica C 418 (2005) 99–106
103
0.30 in situ 0.25
Relative strain (%)
ex situ 0.20
0.15
0.10
0.05
0.00 550
650
750
850
950
Annealing temperature (˚C)
Fig. 3. Development of strain with annealing temperature.
2.50
Density (g/cm3)
2.00
1.50
1.00 doped sample (in situ)
0.50
doped sample (ex situ) pure sample (in situ) pure sample (ex situ)
0.00 550
650
750
850
950
Annealing temperature (˚C)
Fig. 4. Annealing temperature dependence of sample density.
It can be seen from the SEM micrographs in Fig. 5 that both in situ and ex situ samples annealed at 800 °C are porous. The in situ sample consists of larger grains (regions 3 and 5) with particle size ranges from a few lm to as large as about 15 lm and agglomerates of fine grains (are apparent regions 1, 2, 4 and 6). As shown in Table 2, the larger grains contain very low Cu but large amounts of B. Therefore, these regions can be considered as MgB2 and CuB24 rich. For regions with
Fig. 5. Scanning electron micrographs of (a) in situ and (b) ex situ Cu doped samples annealed at 800 °C.
fine grains, the Cu content is substantially higher. Although the amount of B is higher than Mg, this is not as great as found in larger grains. In regions where Mg has a higher proportion than B (region 2), the content of Cu is lower than in those where the amount of Mg is lower than B (regions 1 and 4). In region 6, there is only 3 at% of Cu, however, a significant amount of oxygen was also found there. This shows that the fine grains are either Cu or oxygen rich, suggestive of MgCu2 and MgO.
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Table 2 Results from EDS for the in situ and ex situ Cu doped samples Regions
1
2
3
4
5
6
In situ Mg (at%) B (at%) O (at%) Cu (at%)
6.68 30.68 1.9 60.74
36.47 26.14 2.65 34.74
10.67 74.49 5.06 9.77
3.35 29.87 3.77 63.01
14.52 84.04 0.94 0.5
18.15 61.3 17.54 3
30.19 55.08 12.06 2.67
27.22 58.48 9.46 4.84
31.8 53.94 6.5 7.76
45.45 49.92 3.97 0.66
41.9 38.56 6.82 12.72
27.64 51.38 10.57 10.41
Total (100 at%) Ex situ Mg (at%) B (at%) O (at%) Cu (at%) Total (100 at%)
SEM images in Fig. 5b show that the ex situ sample have randomly distributed fine grains with sizes as large as a few lm and some of them are as small as hundreds of nm. From the EDS results, the Mg content in regions 1–6 is either lower than or comparable to B. This indicates that reaction between MgB2 and Cu is not vigorous, as much of the MgB2 still remained. The oxygen and Cu content in these samples shows no significant difference except in region 4 where Cu content is low. Relatively, the ex situ samples have lower Cu but higher oxygen content than the in situ samples. The higher amount of oxygen is probably due to MgO contamination in the commercial MgB2 powder. 3.3. Superconducting properties Fig. 6 show the results from AC susceptibility measurements in the temperature range of 5– 65 K. The Tc onset which was determined from the deviation point of transition curve is shown in Fig. 7. Tc is higher in the in situ samples than in the ex situ samples. The Tc of both in situ and ex situ samples only shows minor changes within the annealing temperature range of 600 °C– 900 °C. The variation of superconducting transition temperature, DT for in situ and ex situ samples is 0.5 K and 1.2 K, respectively. For in situ samples, the transition is sharp except for the sample annealed at 900 °C in which Tc could not be detected as the diamagnetic signal
is weak (Fig. 6a). There is no significant change of superconducting volume fraction in the annealing temperature range of 600 °C–800 °C. The small variation in vint can be accounted for by calculation error. The large magnitude of susceptibility (<8) indicates the bulk superconductivity in these samples. At 900 °C, the drastic drop of the superconducting volume fraction mainly resulted from a decrease in the phase volume of MgB2 by forming more secondary phases such as MgB4, MgB6 and MgCu2. Tc only showed a slight increase from 38.0 K at 600 °C up to 38.5 °C at 700 °C before it decreased slightly to 38.4 K at 800 °C. As shown in Fig. 6(b), broadening in the transition temperature is observed in the ex situ samples. Besides, the samples annealed at 800 °C–900 °C exhibit a dual-transition indicating the presence of multiple superconducting phases. Tc was increased from 37.5 K at 600 °C up to 38.3 K at 800 °C before it dropped to 37.1 K at 900 °C. This is accompanied by the same evolution of superconducting volume fraction which increased from 600 °C to 800 °C and decreased at 900 °C (jvint800 °Cj > jvint700 °Cj > j vint600 °Cj > jvint900 °Cj).
4. Summary In situ and ex situ Cu doping in MgB2 has been carried out in the temperature range of 600 °C–
S.K. Chen et al. / Physica C 418 (2005) 99–106 0.0
105
38.6 -0.0010
Susceptibility,χ int
-0.1
-0.3
38.2
-0.0025 -0.0030
38.0
-0.0035
600 ˚C
-0.4
700 ˚C -0.5
-0.0040
T c (K)
Susceptibility,χ int
-0.2
38.4
-0.0015 -0.0020
-0.0045 5
15 25 35 Temperature (K)
45
800 ˚C
-0.6
900 ˚C
37.8 37.6 37.4
-0.7
37.2 in situ
-0.8
37.0
ex-situ
-0.9 5
15
(a)
25
35
45
0.0
700 ˚C
900 ˚C Susceptibility,χ int
750 850 Annealing temperature (˚C)
950
800 ˚C
-0.2 -0.3 -0.4 -0.5 -0.6 -0.7 -0.8 5
(b)
650
Fig. 7. Variation of Tc onset with annealing temperature in in situ and ex situ samples. Solid lines are the guides for the eyes.
600 ˚C
-0.1
36.8 550
Temperature (K)
15
25
35
45
For both in situ and ex situ samples, Cu doping leads to an increase of density in sample annealed at 800 °C by 28% and 26%, respectively. SEM micrographs show that the in situ samples consist of larger grains with low Cu content and agglomerates of fine grains corresponds to Cu rich regions. On the other hand, compositional variation in the ex situ sample is less pronounced. The increasing annealing temperature causes only minor changes in the Tc onset but decreases the superconducting volume fraction.
Temperature (K)
Fig. 6. AC susceptibility measured at applied AC field of 1 Oerms and 333.3 Hz for (a) in situ and (b) ex situ Cu doped samples. The inset shows the sample annealed at 900 °C.
900 °C. Attempts to substitute Cu into MgB2 resulted in the formation Mg–Cu phases but showed no evidence of Cu doping into the lattice structure. Within the limits of error, no significant lattice strain nor alteration in a- and c-axes was introduced by varying annealing temperature. This indicates that all the samples suffered only minor Mg deficiency. At higher annealing temperatures (P800 °C), the formation of MgB4 and MgB6 is probably due to a loss of Mg to Mg–Cu alloys.
Acknowledgments We would like to thank R. Stern and D. Smith for the technical assistance. S.K. Chen acknowledges UPM and EPSRC for the financial support.
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