Fabrication of perovskite-type Ba(Sn_1−x Ta _x )O₃ ceramics and their power factors

Abstract

Perovskite-type Ba(Sn_1−x Ta _x )O₃ (0.01 ≤ x ≤ 0.06) ceramics with high relative densities (92.7–94.4 %) were fabricated using the hot isostatic pressing (HIP) method at 1273 K and 196 MPa for 4 h in an atmosphere of argon gas. The lattice parameter decreased slightly with increasing x. From the XPS measurement, the Ta⁵⁺ ion was stable in Ba(Sn_1−x Ta _x )O₃ ceramics and the broad peak of the Ta4f level was the overlap between the Ta⁵⁺4f_5/2 and Ta⁵⁺4f_7/2 levels. Ba(Sn_1−x Ta _x )O₃ ceramics were n-type semiconductors, and their electrical resistivities increased with increasing x. The increase in the electrical resistivity was explained by impurity scattering due to the presence of the Ta ions. The absolute value of the Seebeck coefficient (S) increased with increasing temperature and x. The power factor (S ² σ), which was calculated from electrical conductivity (σ) and the Seebeck coefficient, was ca. 1.0 × 10⁻⁵ W m⁻¹ K⁻² at x = 0.01.

Introduction

BaSnO₃ has a cubic perovskite-type structure with a = 4.116 Å and behaves as an n-type semiconductor with a band gap of 3.4 eV [1, 2]. The Seebeck coefficient (S) of BaSnO₃ was reported to be ca. −300 μV/K at 900 K [3]. Upadhyay et al. measured the electrical conductivity (σ) of BaSnO₃ from 310 to 520 K at 0, 1, 10, and 100 kHz [4]. The dc electrical resistivity (ρ _dc) is 1.6 × 10⁹ Ω cm or more below 350 K and decreases linearly with the rise of temperature. Since the sample was prepared at very high temperature, a slight amount of oxygen was lost from the sample. Since the released electrons were caught by the Sn⁴⁺ ion to generate the Sn²⁺ ion, electrical conduction was explained by the hopping of charge carriers between the Sn⁴⁺ and Sn²⁺ ions. According to Trari et al. [5] (Ba_1−x La _x )SnO₃ (0 ≤ x ≤ 0.02), which was synthesized using a solid-state reaction, is metallic at least for x ≥ 0.0025. Its electrical resistivity is independent of temperature and varies strongly with increasing x. The absolute value of the Seebeck coefficient increases with the rise of temperature and exhibits a large value (ca. −120 μV/K for x = 0.001) at 300 K.

Yasukawa et al. synthesized Ba(Sn_1−x Co _x )O₃ using a solid-state reaction and measured both electrical conductivity and the Seebeck coefficient [6]. The samples have a relative density of ca. 95 % and are p-type semiconductors. Electrical conductivity is 10⁻¹–10² Ω⁻¹ m⁻¹ in the measured temperature range and increases with increasing x. They proposed that p-type conduction occurs by the hopping of holes between the localized states of the Co ions. The Seebeck coefficient increases with the rise of temperature and exhibits a large value (ca. 80 μV/K) at 770 K. The power factor (S ² σ) is frequently used to evaluate the performance of thermoelectric materials. The sample (x = 0.15) has a high power factor (ca. 2.7 × 10⁻⁷ W m⁻¹ K⁻²) at 1073 K.

According to Singh et al. [7], the electrical resistivities of the Ba(Sn_1−x Nb _x )O₃ ceramics that were sintered at high temperatures decreases with increasing x, has a minimum value at x = 0.01, and then increases with increasing x. The decrease in the electrical resistivity in the range of 0.001 ≤ x ≤ 0.01 was explained by the electrons released from both the Nb ion and an oxygen vacancy. On the other hand, the increase in the electrical resistivity in the range of 0.01 ≤ x ≤ 0.10 was explained by impurity scattering due to the presence of the Nb⁵⁺ ions. To improve the power factor of BaSnO₃, electrical conductivity must be increased. In the present study, Ba(Sn_1−x Ta _x )O₃ ceramics with high relative densities were fabricated using the hot isostatic pressing (HIP) method. Electrical conductivity and the Seebeck coefficient were measured to investigate the potential of these oxides as thermoelectric materials.

Experimental

The Ba(Sn_1−x Ta _x )O₃ samples were prepared by a solid-state reaction. High-purity powders of Ba(NO₃)₂ (Rare Metallic, Japan, 99.99 %), SnO₂ (Kojundo Chemical Laboratory, Japan, 99.99 %), and Ta₂O₅ (Rare Metallic, 99.99 %) were weighed in the desired proportions and milled for 1 h in an agate mortar. The mixed powder was calcined at 813 K for 1 h in flowing argon and subsequently was calcined at 853 K for 6 h in flowing argon. The calcined powder was uniaxially pressed into a pellet at 49 MPa. The green compact, which was pressed at 245 MPa using the cold isostatic pressing (CIP) method after being wrapped in platinum, was covered with a boron nitride (BN) powder and was sealed in a Pyrex glass tube under vacuum. Densification was achieved using the hot isostatic pressing (HIP) method for 4 h at 1273 K in an atmosphere of argon gas, and the maximum pressure was 196 MPa [8].

The crystal phases of the sintered samples were examined by powder X-ray diffraction (XRD, RAD-2200, Rigaku, Japan) using monochromatic CuKα radiation (Kα₁ = 1.5405 Å). The lattice parameter was determined from high-angle reflections with Si as a standard. The bulk density was measured using the Archimedes method. Scanning electron microscopy (SEM, Model JSM-7001FD, JEOL, Japan) was used for microstructural observation. The atomic percentages of Ba, Sn, and Ta were determined using energy-dispersive X-ray spectroscopy (EDS, Model JED-2300, JEOL). The binding energy of the Ta 4f level was measured using X-ray photoelectron spectroscopy (XPS, Model AXIS-165, Kratos, UK) with monochromatic MgKα radiation (1253.6 eV) at room temperature. The binding energy was calibrated against the C1 s level from residual hydrocarbon molecules. Electrical conductivity and the Seebeck coefficient were measured using the electrical conductivity & Seebeck coefficient measurement system (RZ2001i, Ozawa, Japan).

Results and discussion

According to Upadhyay et al. [9], the relative density of BaSnO₃ that was sintered at 1525 K in air was ca. 80 %. In the present study, therefore, the HIP method was used to obtain Ba(Sn_1−x Ta _x )O₃ ceramics with high relative densities. Figure 1 shows the XRD patterns of both a mixture (x = 0.06) of the starting materials and the Ba(Sn_0.94Ta_0.06)O₃ (x = 0.06) ceramic fabricated using the HIP method. After HIP processing, all diffraction peaks of the XRD pattern were completely indexed as the perovskite-type structure in the range of 0.01 ≤ x ≤ 0.06 [1]. The lattice parameter decreased slightly from 4.1204 Å (x = 0.01) to 4.1186 Å (x = 0.06), as shown in Table 1. The relative densities of Ba(Sn_1−x Ta _x )O₃ ceramics were high (92.7–94.4 %) regardless of x, as shown in Table 1. SEM images of the fracture surface are shown in Fig. 2. The average particle size decreased slightly from ca. 2 µm (x = 0.01) to ca. 1.5 µm (x = 0.06). The atomic percentage of Ta determined from the EDS measurement increased linearly with increasing x, as shown in Table 1. The results of the XRD and EDS measurements indicated that the Sn ions had been replaced by the Ta ions in Ba(Sn_1−x Ta _x )O₃ ceramics.

Fig. 1

XRD patterns: a a mixture (x = 0.06) of the starting materials, b the Ba(Sn_0.94Ta_0.06)O₃ (x = 0.06) ceramics fabricated using the HIP method at 1273 K and 196 MPa for 4 h in an atmosphere of argon gas

Table 1 The lattice parameter (a), relative density (D), atomic percentage, and activation energy for hopping (E _h) of the Ba(Sn_1−x Ta _x )O₃ ceramics

Fig. 2

SEM images of the fracture surfaces of the Ba(Sn_1−x Ta _x )O₃ (x = 0.01, 0.03, and 0.06) ceramics

The measured Ta4f XPS spectra of the Ba(Sn_1−x Ta _x )O₃ ceramics are shown in Fig. 3. The peak of the Ta4f level was broad and unsymmetrical. Hashimoto et al. used an anodic oxidized Ta₂O₅ thin film on a tantalum plate for the XPS measurement [10]. The binding energies of the Ta⁵⁺, Ta⁴⁺, Ta³⁺, and Ta⁰ levels were 27.5–29.2, 28.6–25.6, 24.0–26.0 and 22.4–24.0 eV, respectively. Yan et al. measured the XPS spectra of LiTaO₃ wafers and reported that the binding energies of the Ta4f level were ca. 28.4 (Ta⁵⁺) and ca. 27.1 eV (Ta⁴⁺) [11]. In the present Ba(Sn_1−x Ta _x )O₃ ceramics, since the binding energy of the Ta4f level was 29.3–29.7 eV, it is thought that the Ta⁵⁺ ion was stable in Ba(Sn_1−x Ta _x )O₃ ceramics and that the broad peak was the overlap between the Ta⁵⁺4f_5/2 and Ta⁵⁺4f_7/2 levels [11]. According to Shannon [12], the ionic radii of the Sn⁴⁺ ion and the Ta⁵⁺ ion with a coordination of 6 are 0.69 Å and 0.64 Å, respectively. Therefore, the slight decrease in the lattice parameter as shown in Table 1 was explained by the difference in the ionic radii between the Sn⁴⁺ and Ta⁵⁺ ions.

Fig. 3

The measured Ta4f XPS spectra of the Ba(Sn_1−x Ta _x )O₃ ceramics

The electrical resistivities (ρ) of the Ba(Sn_1−x Ta _x )O₃ ceramics decreased slightly with the rise of temperature (T), and the log ρ–1000/T curve deviated from a straight line at high temperature, as shown in Fig. 4. Omeiri et al. reported that the electrical resistivity of BaSnO₃ was very high and was assumed to be more than 10⁸ Ω cm [13]. However, the electrical resistivities of the present Ba(Sn_1−x Ta _x )O₃ ceramics were extremely low at x = 0.01 and increased with increasing x in the range of 0.01 ≤ x ≤ 0.06. Since the variation in the relative densities of Ba(Sn_1−x Ta _x )O₃ ceramics was 2 % or less, it is thought that relative density hardly affected the electrical resistivity. According to Singh et al. [7], the electrical resistivity of Ba(Sn_1−x Nb _x )O₃ decreased with increasing x in the range of 0.001 ≤ x ≤ 0.01 and then increased with increasing x in the range of 0.01 ≤ x ≤ 0.10. The decrease in the electrical resistivity up to x = 0.01 was explained by an increase in the number of electrons released from both the Nb ion on the Sn lattice site and an oxygen vacancy. On the other hand, the increase in the electrical resistivity for x > 0.01 was explained by an increase in impurity scattering due to the presence of Nb⁵⁺ ions. In the present Ba(Sn_1−x Ta _x )O₃ ceramics, XPS and EDS measurements indicated that the valence of the Ta ion was 5+ and that the atomic percentage of Ta increased with increasing x. Therefore, it is thought that the electrical resistivities of the Ba(Sn_1−x Ta _x )O₃ ceramics increased as a result of impurity scattering due to the presence of Ta⁵⁺ ions in the same manner as Ba(Sn_1−x Nb _x )O₃ [7].

Fig. 4

The relationship between log ρ of the Ba(Sn_1−x Ta _x )O₃ ceramics and 1000/T

Singh et al. reported that the released electrons in Ba(Sn_1−x Nb _x )O₃ were captured by the Nb⁵⁺ ions or the Sn⁴⁺ ions, and that the Nb⁴⁺ ions or the Sn³⁺ ions were generated [7]. Accordingly, electrical conduction occurred by the hopping of electrons between the Nb⁵⁺ and Nb⁴⁺ ions or between the Sn⁴⁺ and Sn³⁺ ions. Electrical conductivity (σ) for hopping in a small polaron model is given by.

$$ \sigma = \frac{A}{T}{ \exp }\left( { - \frac{{E_{\text{h}} }}{{k_{\text{B}} T}}} \right), $$

where A is a constant, E _h is the activation energy for hopping, and k _B is the Boltzmann constant [14]. Since the log σT–1000/T curves of the Ba(Sn_1−x Ta _x )O₃ ceramics were linear in the measured temperature range as shown in Fig. 5, it is thought that Ba(Sn_1−x Ta _x )O₃ ceramics exhibited electrical conductivity for hopping in the small polaron model. We could not obtain data on electrical resistivity or the Seebeck coefficient above 463 K because the ceramics (x = 0.01) broke during the measurement. According to Taguchi [15], the electrical resistivity of K₂NiF₄-type Ca₂(Mn_1−x Nb _x )O₄ decreased extremely in the range of 0 ≤ x ≤ 0.05 because electron transfer on the Mn³⁺–O–Mn⁴⁺ path became active as the Mn³⁺ ion increased. In the range of 0.05 ≤ x ≤ 0.15, however, the decrease in electrical resistivity was very slight because the Nb⁵⁺ ion, which had no 4d electron, prevented electron transfer on the Mn³⁺–O–Mn⁴⁺ path. In the present Ba(Sn_1−x Ta _x )O₃ ceramics, it is predicted that the Ta⁵⁺ ion, which had no 5d electron, prevented electron transfer on the Sn–O–Sn path. The activation energy for hopping (E _h) increased slightly from 0.02 (x = 0.01) to 0.05 eV (x = 0.06), as shown in Table 1. This result indicates that impurity scattering due to the presence of the Ta ions was more dominant than conduction via hopping in the Ba(Sn_1−x Ta _x )O₃ ceramics.

Fig. 5

The relationship between logged σT of the Ba(Sn_1−x Ta _x )O₃ ceramics and 1000/T

The temperature dependence of the Seebeck coefficient (S) of the Ba(Sn_1−x Ta _x )O₃ ceramics is shown in Fig. 6. Since the Seebeck coefficient was negative, Ba(Sn_1−x Ta _x )O₃ ceramics were n-type semiconductors and the carrier was an electron. The absolute value of the Seebeck coefficient increased with increasing temperature and x. The temperature dependence of the power factor (S ² σ) of the Ba(Sn_1−x Ta _x )O₃ ceramics is shown in Fig. 7. The power factor increased with the rise in temperature and decreased with increasing x despite the slight increase in the absolute value of the Seebeck coefficient. The Ba(Sn_0.99Ta_0.01)O₃ ceramics had electrical conductivity with 35–42 Ω⁻¹ cm⁻¹ and the absolute value of the Seebeck coefficient was 43–53 μV/K. Therefore, the Ba(Sn_0.99Ta_0.01)O₃ ceramics had a maximum power factor (ca. 1.0 × 10⁻⁵ W m⁻¹ K⁻²) at 463 K. Yasukawa et al. reported the thermal properties of the Ba(Sn_1−x Co _x )O₃ ceramics prepared at 1673 K in air using a solid-state reaction [6]. Electrical conductivity and the Seebeck coefficient at 470 K were 0.016–0.06 Ω⁻¹ cm⁻¹ and 30–45 μV/K, respectively. Although there was almost no difference in the absolute value of the Seebeck coefficient between the Ba(Sn_1−x Co _x )O₃ ceramics and the Ba(Sn_1−x Ta _x )O₃ ceramics, the Ba(Sn_1−x Co _x )O₃ ceramics had lower electrical conductivity by two to three orders of magnitude compared with the Ba(Sn_1−x Ta _x )O₃ ceramics. Therefore, the power factor of the Ba(Sn_0.85Co_0.15)O₃ ceramics was very low (ca. 6.0 × 10⁻⁹ W m⁻¹ K⁻² at 470 K). The (Ba_0.95La_0.05)SnO₃ ceramics were synthesized using a polymerized complex method [16]. Electrical conductivity was 0.2–0.4 Ω⁻¹ cm⁻¹ at 400–700 K. The absolute value of the Seebeck coefficient was 40–60 μV/K at 400–700 K; this was almost equal to that of the Ba(Sn_1−x Ta _x )O₃ ceramics. The power factor was ca. 1.0 × 10⁻⁷ W m⁻¹ K⁻² at 400–700 K. Subsequently, Yasukawa et al. fabricated (Ba_1–x La _x )SnO₃ ceramics with relative densities of 93–97 % using the subsequent spark plasma sintering (SPS) technique [3]. The absolute value of the Seebeck coefficient and electrical conductivity at 400–700 K were 40–65 μV/K and 400–500 Ω⁻¹ cm⁻¹, respectively. Although there was no difference in the Seebeck coefficients between the (Ba_1−x La _x )SnO₃ ceramics fabricated using the polymerized complex method and those fabricated using the SPS method, the latter had greater electrical conductivity. As a result, the power factor of the (Ba_0.99La_0.01)SnO₃ ceramics using the SPS method was high (ca. 2.8 × 10⁻⁴ W m⁻¹ K⁻²) at 1073 K. These results indicate that the improvement of the electrical conductivity of the BaSnO₃ ceramics is very important to obtain a high power factor.

Fig. 6

The temperature dependence of the Seebeck coefficient (S) of the Ba(Sn_1−x Ta _x )O₃ ceramics

Fig. 7

The temperature dependence of the power factor (S ² σ) of the Ba(Sn_1−x Ta _x )O₃ ceramics

Conclusions

The perovskite-type Ba(Sn_1−x Ta _x )O₃ ceramics were fabricated in the range of 0.01 ≤ x ≤ 0.06 using the HIP method at 1273 K and 196 MPa for 4 h in an atmosphere of argon gas. The relative densities were 92.7–95.1 % regardless of x. The lattice parameter decreased slightly from 4.1204 Å (x = 0.01) to 4.1186 Å (x = 0.06), and the atomic percentage of Ta increased linearly with increasing x. The electrical resistivity increased with increasing x, indicating that impurity scattering due to the presence of Ta⁵⁺ ions was more dominant than conduction via hopping between the Ta⁵⁺ and Ta⁴⁺ ions or between the Sn⁴⁺ and Sn³⁺ ions. Ba(Sn_1−x Ta _x )O₃ ceramics were n-type semiconductors, and the absolute value of the Seebeck coefficient increased with increasing temperature and x. The power factor was high (ca. 1.0 × 10⁻⁵ W m⁻¹ K⁻²) at x = 0.1 because the decrease in σ exceeded the slight increase in S ².