Thermal and superconducting properties of glass–ceramic HTc BiSrCa(CuPr)O system

j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 1 9 6 ( 2 0 0 8 ) 365–372

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Thermal and superconducting properties of glass–ceramic HTc BiSrCa(CuPr)O system M.A. Aksan ∗ , M.E. Yakinci ¨ ¨ u¨ Universitesi, ¨ ¨ um ¨ u, ¨ Superiletken ¨ I˙non Fen Edebiyat Fakultesi, Fizik Bol Aras¸tırma Laboratuvarı, 44069 Malatya, Turkey

a r t i c l e

i n f o

a b s t r a c t

Article history:

In this study, the glass samples with nominal composition of Bi2 Sr2 Ca2 Cu3−x Prx O10+y , where

Received 23 March 2006

x = 0.5, were fabricated using conventional melt quenching technique. After appropriate heat

Received in revised form

treatments, the glass–ceramic samples obtained showed a deformed and multiphase struc-

13 March 2007

ture. While the a(=b)-axis was nearly unchanged, the c-axis was obtained to be smaller than

Accepted 29 May 2007

that of the unsubstituted Bi-2223 system. From the resistance versus temperature measurements, a semiconducting type behavior down to Tc was observed. The best Tc value was obtained to be 46 K and Tzero 29 K for the sample heat treated at 1123 K (850 ◦ C) for 60 h.

Keywords:

The Pr-substituted Bi-2223 glass–ceramic samples showed negative thermoelectric power,

High-Tc superconductors

indicating that the dominant carriers are electrons in the samples. “Two-band model with

Crystallization kinetics

linear T-term” was used to analyze the TEP data. A rapid rise in thermal conductivity, , was

Thermal conductivity

obtained just below Tc . This suggests an enhancement of the quasiparticle contribution to

Thermoelectric power

the heat conductivity and so an increase of the quasiparticle mean free path. © 2007 Elsevier B.V. All rights reserved.

1.

Introduction

The Bi–O, Sr–O and Ca–O planes in the BSCCO system behave as the charge reservoir blocks (CRB) (Qian et al., 1999). The substitutions on the CRB cause the significant changes in the superconducting properties. However, the substitutions on the Cu sites have the much stronger effect on the superconducting properties due to the scattering of the carriers and/or due to the changes of the carrier concentration on the CuO planes. The substitution of some elements such as Ni, Fe, Co, Zn, V, Sn for the Cu sites caused the suppressed superconducting properties. Intergrain properties of the Bi-based HTc system were influenced by the partial substitution of Cu by 3d-elements, such as Ni, Co and Fe (Pop et al., 1996a, 1996b, 1997; Pop, 1999; Gu et al., 1998). The Tc value decreased significantly with the substitution of Ni and Fe. Especially, in the case of substitution of Fe, the point defects contribute to the

∗

flux-pinning but increase of the point defects on the Cu planes has a destructive effect on the long-range order of Cu interaction (Wang et al., 2000). Co and Zn substitutions result in the suppression of the Tc value stronger than for Ni substitution (Liu et al., 1999; Kuo et al., 1999; Hedt et al., 1994). The depression in Tc was attributed to the Abrikosov and Gor’kov pair-breaking mechanism in the CuO planes. Carrier localization due to the structural disorders induced by substitutions is also useful to explain the decrease in Tc (Doniach and Iniu, 1990). Although in the case of the Sn substitution for Cu in the Bi-2212 system, superconductivity at the low substitution levels was obtained, superconductivity was strongly suppressed at the high substitution levels, which was attributed to the hole filling mechanism (Li et al., 2000a, 2000b, 2000c). However, the suppression of the superconductivity in the V-substituted Bi-2223 system was ascribed to a decrease of the carrier concentration (Nkum, 1998).

Corresponding author. Tel.: +90 422 3410010x3720; fax: +90 422 3410037. E-mail address: [email protected] (M.A. Aksan). 0924-0136/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2007.05.042

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Many studies have been done on the Pr-substitution for Ca in the Bi-based HTc system (Sun et al., 1998a, 1998b; Chen et al., 1994; Li et al., 2000a, 2000b, 2000c; dos Santos et al., 2001). However, there is no study on the Pr-substitution for Cu in the BSCCO HTc materials in literature. In this work, we have prepared the Bi2 Sr2 Ca2 Cu3−x Prx O10+y , where x = 0.5, glass–ceramic samples and investigated the structural, thermal and transport properties of the system obtained.

2.

Experimental details

Glass samples with nominal composition of Bi2 Sr2 Ca2 Cu3−x Prx O10+y , where x = 0.5, was prepared using the glass–ceramic technique. Firstly, an appropriate amount of high purity (at least 99.9%) Bi2 O3 , SrCO3 , CaCO3 CuO and Pr6 O11 oxide powders was mixed in an agate mortar for 3–5 h. The mixture was melted in an ␣-alumina crucible in high temperature melting furnace at 1523 K (1250 ◦ C) for 3–4 h and then the molten material was quenched between two cold copper plates. Thus, rapidly quenched, dark, shiny and approximately 1–3 mm thick amorphous materials were obtained. The heat treatments of glass samples were done in P.I.D. controlled tube furnace in oxygen atmosphere at 1123 K (850 ◦ C) for 40–120 h, considering differential thermal analysis (DTA) results. The superconductivity in the HTc systems is very sensitive to the oxygen content in the unit cell. The oxygen deficiencies in the unit cell which may form during preparation leads to suppression in the superconductivity. In order to compensate these oxygen deficiencies, the materials prepared should be sintered under the oxygen atmosphere. After x = 0.5 substitution level, no glass formation was obtained even at very high temperature range, 1323–1673 K (1050–1400 ◦ C), therefore, those samples were not investigated. Superconducting phases and structural changes during crystallization were analyzed by using X-ray diffraction (XRD). The XRD diffraction data were taken by using a Rigaku RadB system with Cu K␣ radiation and between 3◦ and 60◦ in continuous mode. The scan rate was selected as 3◦ /min.

Differential thermal analysis (DTA) is carried out to determine the nature of the phase transformations within the material prepared, the optimum conditions for heat treatment cycles of the glass ceramic materials and to investigate reaction kinetics of materials. For this purpose, 20 mg crushed and grounded glass samples was examined using Shimadzu system 60 thermal analysis network based on non-isothermal kinetic theory with ␣-alumina reference material. In the crystallization kinetic studies, more than three points are necessary for accuracy of the results. Therefore, four different uniform heating rates were selected for the kinetic studies. Resistance versus temperature of the annealed samples was carried out using four probe dc resistance measurement system with closed cycle He refrigerator (Leybold LT 10 system). Highly conductive silver paint was used for contacts. Thermoelectric power (TEP) and thermal conductivity were measured as a function of temperature with a steady-state heat-flow technique using closed cycle He refrigerator in the temperature range of 10–300 K. Temperature gradient across the sample with 1–2 K/cm was monitored. Thermoelectric voltages were recorded using Keithley 182 nanovoltmeter and the temperature gradient using two silicon diodes. At any temperature, the voltage difference (V) was recorded to compute the TEP (S = V/T). All these steps were done automatically as in the resistance–temperature measurement. The analysis of the thermoelectric power data was done by using “Two-band model with linear terms” (Forro et al., 1989, 1990).

3.

Result and discussion

3.1.

XRD results

The XRD pattern of the melt-quenched x = 0.5 Pr-substituted sample was shown in Fig. 1a. A broad halo around 2 = 30◦ was obtained, which is the characteristic of the glass materials and indicates the absence of any atomic arrangement and periodicity in three dimensions.

Fig. 1 – X-ray diffraction patterns of the Bi2 Sr2 Ca2 Cu2.5 Pr0.5 O10+y system; (a) glass sample, (b) the sample heat treated at 1123 K (850 ◦ C) for 40 h, (c) the sample heat treated at 1123 K (850 ◦ C) for 60 h and (d) the sample heat treated at 1123 K (850 ◦ C) for 120 h.

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Table 1 – Crystallographic and electrical properties of the Bi2 Sr2 Ca2 Cu2.5 Pr0.5 O10+y system Heat treatment ◦

1123 K (850 C)–40 h 1123 K (850 ◦ C)–60 h 1123 K (850 ◦ C)–120 h

Atmosphere

˚ Crystalline size (A)

Tc (K)

Tzero (K)

T (K)

Hole concentration p

O2 O2 O2

258 272 295

43 46 37

20 29 22

23 17 15

0.238 0.243 0.249

Tc is superconducting transition temperature, Tzero the zero resistance temperature, T = Tc − Tzero and p is the hole concentration per CuO planes.

XRD patterns of the samples are shown in Fig. 1b–d. The Bi-2212 phase was identified to be main phase in all the samples prepared. However, Bi-2201 and the impurity phases such as CuPr2 O4 , CuPrO2 , Cu2 SrO3 and Bi1.35 Pr0.65 O3 were obtained. The intensity of the peaks of the superconducting phases increased with increasing the heat treatment time, suggesting that the crystallization of the superconducting phases increased with increasing the heat treatment time, but no change in the impurity phase formation was found. This obviously indicates that extension of the heat treatment time has not any effect on the thermodynamic activities of the impurity ions in the Pr-substituted samples. When the crystalline size of the samples was computed by Scherrer Method (Cullity, 1978) given by L=

0.9 t cos 

(1)

where L is crystalline size, t FWHM of the XRD peaks,  the peak ˚ angle and  is the wavelengths of the X-rays used ( = 1.5405 A in our system). The results are listed in Table 1. The aver˚ for the sample heat age crystalline size increased from 258 A ˚ for the sample heat treated at 1123 K (850 ◦ C) for 40 h to 295 A ◦ treated at 1123 K (850 C) for 120 h. The reduction in superconducting transition width, Tc (Table 1) can be attributed to the increase in the crystalline size by increasing the sintering time. The phase coordination was destroyed in the Pr-substituted samples and the multiphase, complex and deformed structure was obtained. Crystal symmetry was found to be tetragonal ˚ and the unit cell parameters were calculated as a(=b) = 5.3990 A ˚ and c = 31.3011 A. While the a(=b)-axis was nearly unchanged, the c-axis decreased with the substitution of Pr, compared ˚ and c = 37.110 A). ˚ The to the Bi-2223 system (a(=b) = 5.4012 A change of the unit cell parameters is attributed to the excess oxygen arising from the replacement of Cu2+ by Pr3+/4+ and the sintering under O2 atmosphere. The excess oxygen atoms incorporate into the Bi–O layers in the BSCCO system and these excess oxygen atoms on the Bi–O layers cause a decrease

Fig. 2 – DTA curves of the Bi2 Sr2 Ca2 Cu2.5 Pr0.5 O10+y glass system; heating rates: (a) ˛ = 278 K/min, (b) ˛ = 283 K/min, (c) ˛ = 293 K/min and (d) ˛ = 303 K/min.

in the net positive charge. Therefore, the repulsion between the Bi–O layers decreases and this leads to the shrinkage of the SrO–BiO–BiO–SrO slabs, i.e. a decrease in the c-axis. In addition, as far as the ionic sizes of both Pr (3+ and/or 4+) and Cu (2+) are concerned, the ionic radii of Pr3+ and/or Pr4+ are 1.013 ˚ According to the decrease ˚ while for Cu2+ it is 0.72 A. and 0.90 A, ˚ we can conclude that Pr in the c-axis from 37.110 to 31.3011 A, exists between 3+ and 4+ valence states in the system. Similar results were obtained by Sun et al. in the Pr-substituted Bi-2212 system for Ca (Sun et al., 1998a, 1998b).

3.2.

DTA results

The DTA curves of the x = 0.5 Pr-substituted Bi-2223 sample with the heating rates of ˛ = 278, 283, 293 and 303 K/min were shown in Fig. 2a–d, respectively, and the DTA data in Table 2. The glass transition temperature, Tg , marked by an arrow

Table 2 – The DTA data of the x = 0.5 Pr-substituted Bi-2223 system Heating rate (K/min) 278 283 293 303

Tg (K)

Tx1 (K)

T (Tx1 − Tg ) (K)

Tx2 (K)

Tendo (K)

Tpm (K)

728 730 731 733

735 739 742 757

7 9 11 24

1027 1044 1035 1051

1165 1173 1169 1174

1197 1198 1199 1205

Tg is the glass transition temperature, Tx1 the first crystallization temperature, T (Tx1 − Tg ) the glass working range, Tx2 the second crystallization temperature, Tendo the temperature of the endotermic peak and Tpm is the partial melting temperature as explained in Section 3.2.

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Fig. 3 – The Augis–Bennett and Kissinger plots of the x = 0.5 Pr-substituted system. The rounded symbols shows the points calculated based on the Augis–Bennet model and the square symbols the points calculated based on the Kissinger model.

in Fig. 2, was found to be between 728 and 733 K. The first crystallization temperatures, Tx1 , were obtained between 735 and 757 K, depending on the heating rate. We believed that these exothermic activities correspond to temperatures at which the impurity phases such as CuPr2 O4 , CuPrO2 , CuSrO3 and Bi0.35 Pr0.65 O3 detected in XRD results begin to grow. The second crystallization temperatures, Tx2 , between 1027 and 1051 K were obtained, which are attributed to the formation of the Bi-2201 phase. After 1123 K, two endothermic peaks were observed; the first peak between 1165 and 1174 K, Tendo , shows the formation temperature of the Bi-2212 phase and the second between 1197 and 1205 K the partial melting temperature, Tpm , of the sample, at which the sample starts to melt. Important results can be derived from the DTA data; the crystallization temperatures, Tx , were shifted towards the higher temperatures with increasing the heating rate. Increase at Tx depends on the nucleation rate in the material. When the heating rate is high, the material stays in the temperature region of nucleation for a time shorter than time lag. Vice versa when the heating rate is low, the samples stay in the temperature region of nucleation for times longer than time lag and the formation of secondary nuclei might take place (Karamanov and Pelino, 2001). In addition, increase of Tx in the system reveals the increase in the surface nucleation in the materials (Aksan et al., 2005). The activation energy for crystallization of the system prepared was computed by using Kissinger (Kissinger, 1956; Sunol et al., 2003) and Augis–Bennett method (Augis and Bennett, 1978). Fig. 3 shows Kissinger and Augis–Bennett plots and the results obtained were listed in Table 3. The obtained activa-

Table 3 – The activation energy for crystallization of the unsubstituted (Balcı, 1997) and the x = 0.5 Pr-substituted samples X-value 0.0 0.5

Augis–Bennett model (kJ/mol)

Kissinger model (kJ/mol)

339.18 377.37

336.46 374.75

Fig. 4 – Avrami parameter, n, for the x = 0.5 Pr-substituted samples obtained by Eq. (1).

tion energies of the Pr-substituted system are higher than that of the unsubstituted Bi-2223 system (Balcı, 1997) by almost 38 kJ/mol (see Table 3). The crystallization of a glass system needs the rearrangement of unlike atoms, during which an atom must overcome the bonding energies with neighbors to take its lattice position of primary crystals. This indicates that effective activation energy for crystallization of a glass sample reflects the interaction of atoms. Therefore, the sample including the atoms which show much stronger attractive interaction produces higher activation energy. Based on the discussion above, the sample prepared in this study shows much higher stability against crystallization than the unsubstituted Bi-2223 system. In addition, increase of the activation energy indicates increase of the surface crystallization in the system and crystals grow at higher viscosity of supercooled melt (Karamanov and Pelino, 2001). The Ozawa equation (Ozawa, 1971; Rao et al., 2004; Sunol et al., 1999), which provides important information about the crystal morphology in the materials prepared, is given by d[ln(− ln(1 − )] = −n d(ln ˛)

(2)

where n is the Avrami parameter, ˛ the heating rate and  is determined at a fixed temperature, T, on the exotherms obtained at different heating rates. The value of  is calculated from the ratio of partial area of the crystallization peak at any temperature to the total area of the crystallization peak. The plot of ln(−ln(1 − )) versus ln ˛ is the straight line whose slope gives the parameter n, Fig. 4. The Avrami Parameter, n, was found to be 2.39. This value indicates that the growth is diffusion-controlled and simultaneous nucleation and three-dimensional parabolic growth takes place in the system during the first crystallization process (Karamanov and Pelino, 2001). For the diffusion-controlled growth, one may have the following cases: 1.5 < n < 2.5 reflects growth of small particles with a decreasing nucleation rate. The higher activation energy for crystallization demonstrates that atomic diffusion in sample is difficult. The difficult atomic diffusion retards the nucleation and growth, resulting in a decreasing nucleation rate. However, the n values higher than 2.5 indicate the growth of small particles with an increasing nucleation rate (Christian, 1975; Yan et al., 2004). Increase of the nucleation rate results in the composition fluctuations in the regions adjacent to the growing nuclei, which cause a

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Fig. 5 – Resistance vs. temperature (R–T) curve of (a) the sample heat treated at 1123 K (850 ◦ C) for 40 h, (b) the sample heat treated at 1123 K (850 ◦ C) for 60 h and (c) the sample heat treated at 1123 K (850 ◦ C) for 120 h.

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Fig. 6 – Temperature dependence of S(T) for Bi2 Sr2 Ca2 Cu2.5 Pr0.5 O10+y system: (a) the sample heat treated at 1123 K (850 ◦ C) for 40 h, (b) the sample heat treated at 1123 K (850 ◦ C) for 60 h and (c) the sample heat treated at 1123 K (850 ◦ C) for 120 h. Solid lines are fitting curves of two-band model with linear T-term to TEP data.

chain reaction-like process leading to an increasing nucleation rate.

3.3.

Resistance versus temperature (R–T) results

The temperature dependence of the resistance of the x = 0.5 Pr-substituted system is shown in Fig. 5a–c. The sample heat treated at 1123 K (850 ◦ C) for 40 h displayed a semiconducting trend down to Tc . After the Tc value was obtained at 43 K, the zero resistance value, Tzero , was reached at 20 K, Fig. 5a. Although the room temperature resistance of the sample heat treated at 1123 K (850 ◦ C) for 60 h is smaller than that of the sample heat treated at 1123 K (850 ◦ C) for 40 h, semiconducting behavior is observed again, as shown in Fig. 5b. Tc and Tzero were found to be 46 and 29 K, respectively. When the heat treatment was done at 1123 K (850 ◦ C) for 120 h, the room temperature resistance increased compared to the previous sample. Semiconducting behavior was obtained and the sample showed a Tc value of 37 K and the Tzero value of 22 K. The observed metal–semiconductor transition clearly indicates a doping-induced localization of charge carriers due to the presence of structural disorder. Decrease of the normal state resistance with increasing the heat treatment time (from 40 to 60 h) exhibits slight reduction in the localization of charge carriers, but the localization does not completely disappear. The growth of the impurity phases at the grain interface in the samples supports the assumption above. In addition, the valence of the substitution element Pr (3+ and/or 4+) destroys the electronic configuration of the samples. These reasons are responsible for the suppressed electrical properties and lower Tc values (<55 K) of the samples. The hole number per the CuO planes was computed by the equation Tc 2 = 1 − 82.6(p − 0.16) Tcmax

(3)

where Tcmax is taken 110 K for the Bi-2223 system and p is the number of the holes per Cu (Presland et al., 1991) and the

obtained results is given in Table 1. Presland et al. was found that the average value of p ranges from 0.16 to 0.116 for Bi2223, from 0.22 to 0.16 for Bi-2212 and higher than 0.22 for Bi-2201. In this study, the hole concentration increased by the substitution of Pr and there is a transition from the hole concentration regime of Bi-2223 (optimal-doped-region) to that of Bi-2201 (overdoped-region). This indicates a reduction in the conduction and suppression of superconductivity. We believed that since the cation state of Pr can be between 3+ and 4+, which is bigger than that of the Cu2+ , extra charges is transferred to the system by the substitution of Pr. In this case, the Cu–O planes are doped with highly hole concentration and so material lies in highly suppressed superconducting and/or non-superconducting domain.

3.4.

Thermoelectric power results

Like electrical conductivity, thermoelectric power of high-Tc materials also gives very important information about the nature the charge carriers. Therefore, thermoelectric power measurement is a useful tool for the characterization of highTc materials. Fig. 6 shows thermoelectric power versus temperature curve (S–T) of the samples prepared. The similar trend was observed in all the samples. The sign of TEP of the sample was obtained to be negative, indicating that electrons were dominant charge carriers in the sample prepared (n-type conductivity). When the temperature was decreased from room temperature (300 K), the non-linear thermal variation in TEP shows a negative peak at almost 55 K which is just above Tc and then falls sharply to zero. The negative peak suppressed with increasing the heat treatment time. Reduction of phonon heat current leads to a reduction in momentum transferred to electrons, which means phonon-drag effect. Therefore, the negative peak in the system is ascribed to phonon-drag effect. The formation of the impurity bonds such as Pr–Cu–O and Pr–O–Bi as identified in XRD investigations, which is the dominant scattering centers in the samples prepared, may be

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Table 4 – The fit parameters A, B, ˛ and estimated values of (E0 − EF ) and  of the system Heat treatment condition ◦

1123 K (850 C) for 40 h 1123 K (850 ◦ C) for 60 h 1123 K (850 ◦ C) for 120 h

A (mV) −1.6712 −1.4013 −0.9320

B (K)

˛ (mV/K2 )

E0 − EF (K)

 (K)

64.782 90.058 113.750

1.53 × 10 1.26 × 105 8.06 × 106

−0.969 −0.813 −0.541

7608.126 14703.260 23456.97

responsible for the non-linear behavior of TEP. Similar results were obtained in some glass–ceramic Bi-based systems by research groups (Chatterjee et al., 1998, 1997; Bhattacharya et al., 1998; Barik et al., 1999). It is worth noting that TEP increased rapidly to zero with further decrease of temperature. This behavior is consistent with the fact that S goes to zero in superconducting state. The TEP data of the samples prepared in the glass–ceramic form were analyzed by using “Two-band model with linear Tterm”.

3.4.1.

Two-band model with linear T-term

Gottwick et al. analyzed the TEP data of CeNi samples assuming a Lorenzian resonance near the Fermi level (Gottwick et al., 1985). In order to analyze the data they used the formulas given by: S=

AT B2 + T 2

A=2

(4)

E0 − EF e

(5) 2

B2 = 3

(E0 − EF ) +  2 2 k2B

(6)

where AT is the conduction of the metallic hole and B/T that of the semiconductor type electrons. E0 and  are the center and the width of the resonance, respectively. The theory is based on a localized band in density of states near the Fermi level, which is superimposed on a broad band. This resonance peak gives the characteristic temperature dependence of S. Forro et al. indicated that the temperature dependence of S of HTc systems is similar to that of mixed valent heavy fermion systems (Forro et al., 1989, 1990). They added a linear term to Eq. (4) to explain the temperature dependence of S in the Bi-2212 single crystals and obtained the following equation: S=

AT + ˛T B2 + T 2

5

Table 4. This is in contradiction to the expected result. The normal band contribution, ˛T, also increases the heat treatment time.

3.5.

Thermal conductivity results (T)

Fig. 7 shows the thermal conductivity–temperature curve (–T) of the Pr-substituted Bi-2223 glass–ceramic material heat treated at 1123 K (850 ◦ C) for 40–120 h. (T) of the samples displayed all the features typical to those known HTc systems. The temperature dependence of (T) was obtained to be linear down to Tc . After a small minimum, a sharp rise with a maximum and then a sharp drop were obtained just below Tc . Such a feature has been observed in most substituted and/or unsubstituted HTc systems (Chanda and Dey, 1994; Bougrine et al., 1998, 2000; Aksan et al., 1999; Yakıncı, 1997). The origin of this small minimum near Tc is not clear but may be attributed to superconducting fluctuation contribution (Castellazi et al., 1997; Houssa et al., 1996; Peacor et al., 1991; Yu et al., 1992). The origin of the sharp maximum has also not been clarified completely. It ascribed to either an electronic contribution or caused by the reduction of the phonon scattering mechanism or both. Recent investigations displayed that sharp rise of (T) gives for an enhancement of the quasiparticle contribution to the heat conductivity and so an increase of the quasiparticle mean free path (Houssa et al., 1996; Houssa and Ausloss, 1995, 1996; Ausloss and Houssa, 1999; Yu et al., 1994; Zeini et al., 2001). It was observed that the magnitude of the thermal conductivity, (T), was strongly influenced by the heat treatment time. Smaller (T) was obtained in low heat treatment time however the magnitude of (T) increased with increasing the heat treatment time. The strong impurity concentration which is

(7)

where ˛T represents normal band contribution. In this model electrons and holes have different mobilities. Eq. (7) was fitted the TEP data obtained by some research groups very well (Sita and Singh, 1998; Wang et al., 1996; Mandal et al., 1992; Singh and Sita, 1999). The fit of Eq. (7) to our experimental data is given with solid line in Fig. 6 and the fitting parameters in Table 4. Experimental points of the samples fit well with Eq. (7). A systematic variation was obtained in the parameters with increase of heat treatment time. (E0 − EF ) and  values increase by increasing the heat treatment time. This suggests that Fermi level should go down relatively to the top of the band with increasing hole density in the system, while (E0 − EF ) value increases as seen in

Fig. 7 – –T curve of the (a) the sample heat treated at 1123 K (850 ◦ C) for 120 h, (b) the sample heat treated at 1123 K (850 ◦ C) for 60 h and (c) the sample heat treated at 1123 K (850 ◦ C) for 40 h.

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detected in XRD investigations reduces the heat transport due to a strong increase of the electron–phonon-impurity scattering. Subgrain boundaries and inhomogeneous oxygen content in the sample has large effect on the intensity of the (T) peak (Knizek et al., 1998).

3.6.

Conclusion

In literature, the effect of Pr-substitution for Ca in the BSCCO HTc system on the superconducting properties was usually investigated but not the Pr-substitution for Cu. In this study, we examined the superconducting, structural and thermal properties in the Pr-substituted glass–ceramic Bi2223 system. The samples with nominal composition of Bi2 Sr2 Ca2 Cu3−x Prx O10+y (x = 0.5) were prepared by conventional glass–ceramic technique. Since we could not obtain any glass materials for higher substitution levels than x = 0.5, the samples after x = 0.5 were not investigated. From the XRD investigations it is seen that the Pr-substitution leads to multiphase, complex structure and impurity phase formation. It suggests that the Pr-substituted material shows high stability against crystallization which is supported by the crystallization kinetics investigations. Resistance–temperature measurements show semiconducting type behavior down to Tc in the system. The Tc and Tzero values decreased in the Pr-substituted system, compared to the pure Bi-2223 system. The doping-induced localization, increase of hole concentration and the growth of the secondary (impurity) phases at grain interface causes a reduction in the conduction and the Tc and Tzero values. The Pr-substituted Bi-2223 system was found to have a negative thermoelectric power, indicating that major charge carriers are electrons in the system. “Two-band model with linear Tterm” was used to analyze the TEP data. The rapid rise of  just below Tc indicates the enhancement of the quasiparticle contribution to the heat conductivity and so increase of the quasi-particle mean free path. It supports an increase on the electrical contribution to the total thermal conductivity below Tc .

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