Effect of A-Site Cation Deficiency on the Thermoelectric Performance

Article pubs.acs.org/JPCC

Effect of A‑Site Cation Deficiency on the Thermoelectric Performance of Donor-Substituted Strontium Titanate A. V. Kovalevsky,*,† A. A. Yaremchenko,† S. Populoh,‡ A. Weidenkaff,‡ and J. R. Frade† †

Department of Materials and Ceramic Engineering, CICECO, University of Aveiro, 3810-193 Aveiro, Portugal Empa, Solid State Chemistry and Catalysis, Ueberlandstr. 129, CH-8600 Duebendorf, Switzerland

‡

S Supporting Information *

ABSTRACT: Donor-substituted strontium titanate is known as one of the best highly performing n-type oxide thermoelectrics. In the present work, structural, microstructural, and thermoelectric properties of Sr1−xPrxTiO3±δ, Sr1−1.5xPrxTiO3±δ, and Sr1−yTi0.8Nb0.2O3±δ (x = 0.05−0.30, y = 0−0.10) perovskite-type titanates were assessed to identify the impact of nominal A-site deficiency on thermoelectric performance. A large increase in power factor was observed for A-site nonstoichiometric materials at high donor-substitution level, provided by the favorable changes in electronic structure, defect chemistry, and microstructure. Complex dependence of the total conductivity and Seebeck coefficient on strontium deficiency suggests effects of microstructure on the electrical properties of Prand Nb-substituted ceramic samples. Formation of oxygen vacancies was found to suppress lattice thermal conductivity at low and intermediate temperatures, whereas the electronic part of the heat transfer increases for cation-deficient materials. For both Pr- and Nbsubstituted titanates, introducing nominal A-site deficiency represents a promising strategy for improving thermoelectric performance.

1. INTRODUCTION Thermoelectric materials, which can convert thermal energy into electric energy (Seebeck effect) and vise versa (Peltier effect) are expected to provide enhanced overall efficiency in energy consumption and prospects for waste heat recovery. Thermoelectric conversion combines simplicity without employing any moving parts, silent operation, excellent scalability, and reliability.1,2 The gap between the actual efficiency of thermoelectric generation and its upper limit (i.e., Carnot efficiency) is determined by the dimensionless figure of merit ZT (σα2 × T/κ), which increases with electrical conductivity (σ) and Seebeck coefficient (α), and decreases with thermal conductivity (κ). Unfortunately, these properties are usually interrelated, with emphasis on simultaneous contributions of charge carriers to electrical and thermal conductivity, and also because high values of Seebeck coefficient are usually found for low carrier concentrations. Major efforts are thus dedicated to seeking materials with the highest mobility of charge carriers combined with other strategies for suppressing lattice contributions to thermal conductivity. The most studied and well-known thermoelectric materials are Bi- and Pb-based chalcogenides (Bi2Te3, Bi2Se3, PbTe-based, etc.), and include some of the best performing thermoelectrics with the figure of merit higher than 1 (ref 3 and references therein). However, these materials are not attractive for practical applications at high temperatures because of the decomposition, vaporization, and melting of the constituents. After discovery of promising thermoelectric performance in layered NaCo2O44,5 and other Co-containing oxide thermoelectrics (ref 6 and references therein), particular interest has © 2014 American Chemical Society

been given to other transition-metal-based mixed-conducting oxide materials with reasonable thermoelectric properties6−12 for absence of toxicity and high natural abundance of constituent elements. Another key issue of thermoelectric technologies is the requirement for both n-type and p-type thermoelectrics. Donor-substituted strontium titanate is one of the most promising n-type oxide thermoelectrics to seek high power factor (σα2)13 based on a specific electronic structure and prevailing lattice contribution to thermal conductivity, which can be tuned by substitution and/or micro/nanoengineering.14−17 Other advantages of SrTiO3-based materials include excellent thermal and phase stability, under both oxidizing and reducing conditions. However, appropriate electrical properties can be attained only after high-temperature heat treatment under very reducing conditions, when donor substitution effectively results in formation of electronic defects. Promising ZT values have thus been reported for reduced bulk polycrystalline titanates with perovskite-type structure, such as ZT = 0.22 for Sr0.9Dy0.1TiO3 at 573 K,18 0.28 for Sr0.9Nd0.1TiO3 at 873 K,19 0.28 for Sr0.88La0.12TiO3 at 773 K,20 0.34 for Sr0.90Pr0.10TiO3 at 1170 K,21 0.35 for SrTi0.8Nb0.2O3 at 1000 K,22 0.36 for La0.1Sr0.83Dy0.07TiO3 at 1045 K,16 and 0.4 at 1040 K for EuTi0.98Nb0.02O3.23 The complex defect chemistry of SrTiO3-based materials is largely determined by oxygen chemical potential in the Received: October 4, 2013 Revised: February 12, 2014 Published: February 12, 2014 4596

dx.doi.org/10.1021/jp409872e | J. Phys. Chem. C 2014, 118, 4596−4606

The Journal of Physical Chemistry C

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Figure 1. Typical room-temperature XRD patterns of Pr- and Nb-substituted titanates (A). The results of refinement for the samples with highest praseodymium content are shown (B, C); the arrow indicates additional peak splitting characteristic for tetragonal-to-orthorhombic transition.

99.9%). Before being weighed, titanium and niobium oxides were annealed at 973 K for 2 h in air. Precursor powders were mixed in the required proportions and consecutively calcined at 1173, 1373, 1473, and 1573 K for 5 h at each temperature, with intermediate regrindings to promote solid-state reactions and to enhance homogenization. The powders were ball-milled with ethanol, and disk-shaped ceramic samples were prepared by uniaxial compacting, followed by sintering. Samples of all compositions, except Sr0.95Pr0.05TiO3±δ, were sintered directly under reducing conditions in flowing 10% H2−90% N2 mixture at 1773 K for 10 h. For Sr0.95Pr0.05TiO3±δ, the upper temperature limit of available furnace with reducing atmosphere was insufficient to prepare the ceramic samples with reasonable densities, i.e., higher than 60%. Therefore, the samples of this composition were preliminary sintered in air for 10 h at 1973 K, to eliminate open porosity, and then reduced at 1773 K for 10 h in 10% H2−90% N2 atmosphere. In the course of sintering or annealing steps, Pt foil or a thick layer of the powder with the same composition was used as substrate in air or hydrogenbased atmospheres, respectively, to avoid reaction with alumina supports. The obtained disk samples were cut into rectangular bars (∼2 × 3 × 12 mm3) for the total conductivity and Seebeck coefficient measurements or polished to provide uniform thickness (∼1.00 mm) for the thermal diffusivity studies. Xray diffraction (XRD), thermogravimetry (TGA), and specific heat capacity studies were performed on the powders prepared by grinding ceramic samples in a mortar. SEM analysis was performed on fracture surface of as-prepared ceramic samples. To assess the compositional homogeneity on the microscale level, the samples were polished, thermally etched, and

atmosphere, whereas dependence on A:B site ratio is still somewhat controversial.24−26 Though effects of A-site cation deficiency have been addressed for potential applications of SrTiO3 with donor additives (e.g., SOFC anode materials24), to the best of our knowledge, the impact on thermoelectric properties of donor-substituted strontium titanates has not been studied. Besides expected changes in defect chemistry, strontium deficiency might be favorable in terms of improved stability if considering possible application in CO2-rich atmospheres at intermediate temperatures. In the present work, Pr- and Nb-substituted titanates were studied as A- and B-site donor-substituted model systems to assess the effect of A-site nonstoichiometry on thermoelectric properties. These materials were selected taking into account known high thermoelectric performance found for Sr(Ti,Nb)O3 oxides, while relevant preliminary investigations of praseodymiumcontaining strontium titanate showed one of the best ZT values among those found for A-site donor-substituted titanates.21 Therefore, the series of studied ceramic materials included Sr1−xPrxTiO3±δ with nominal stoichiometric A:B site ratio, Asite deficient materials with nominal charge compensation for the donor additive Sr1−1.5xPrxTiO3±δ, and A-site deficient materials with partial compensation for the donor additive Sr1−yTi0.8Nb0.2O3±δ, with y = 0, 0.03, 0.06, 0.10.

2. EXPERIMENTAL SECTION Precursor powders of Sr1−xPrxTiO3±δ, Sr1−1.5xPrxTiO3±δ, and Sr1−yTi0.8Nb0.2O3±δ (x = 0.05, 0.10, 0.20, 0.30; y = 0, 0.03, 0.06, 0.10) were prepared by conventional solid-state route using SrCO3 (Sigma Aldrich, ≥ 99.9%), Pr6O11 (Sigma Aldrich, 99.9%), TiO2 (Sigma Aldrich, 99.8%), and Nb2O5 (Aldrich, 4597



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Table 1. Structural Parameters and Density of Sr1−xPrxTiO3±δ, Sr1‑1.5xPrxTiO3±δ, and Sr1−yTi0.8Nb0.2O3±δ ceramics chemical composition

a

structure

space group

lattice parameters (Å)

pseudocubic unit cell volume (Å3)

density (g/cm3)

relative densitya (%)

a = 3.9086(3) a =3.9056(2) a = 5.5294(5) c = 7.8124(6) a = 5.5266(6) c = 7.8061(9) a = 5.5190(3) c = 7.8363(4) a = 5.5108(3) c = 7.8250(4) a = 5.5378(2) b = 7.7971(3) c = 5.5217(2) a = 5.5002(2) c = 7.8122(3) 3.9357(2) 3.9281(1) 3.9287(1) 3.9316(1)

59.71 59.58 59.72

4.85 4.86 5.15

93.7 94.8 98.0

59.61

4.87

94.8

59.67

4.14

76.6

59.41

5.12

98.8

59.61

3.69

66.7

59.09

4.68

89.4

60.96 60.61 60.64 60.77

3.73 5.08 4.82 4.68

71.3 97.7 94.0 93.1

Sr0.95Pr0.05TiO3±δ Sr0.925Pr0.05TiO3±δ Sr0.90Pr0.10TiO3±δ

cubic cubic tetragonal

Pm3m ̅ Pm3m ̅ I4/mcm

Sr0.85Pr0.10TiO3±δ

tetragonal

I4/mcm


tetragonal

I4/mcm


tetragonal

I4/mcm


orthorhombic

Imma


tetragonal

I4/mcm

SrTi0.8Nb0.2O3±δ Sr0.97Ti0.8Nb0.2O3±δ Sr0.94Ti0.8Nb0.2O3±δ Sr0.90Ti0.8Nb0.2O3±δ

cubic cubic cubic cubic

Pm3m ̅ Pm3m ̅ Pm3m ̅ Pm3m ̅

Calculated assuming that δ = 0.

values of thermal conductivity was less than 10% for all measured samples.

examined by scanning electron microscopy and energy dispersive spectroscopy (SEM/EDS). X-ray diffraction patterns were recorded using a Rigaku D/ Max-B diffractometer (Cu Kα; 2Θ = 10−80°; step, 0.02°; exposition, 3 s). Unit cell parameters were calculated from the diffraction data using the profile-matching method in Fullprof software.27 Microstructural characterization was performed by scanning electron microscopy (SEM, Hitachi S-4100 and SU-70 instruments) coupled with energy dispersive spectroscopy (EDS, Rontec UHV and Bruker Quantax 400 detectors, respectively). The experimental densities (ρ) of disk-shaped ceramics were determined by geometrical measurements and weighing. Thermogravimetric analysis (TGA, Setaram SetSys 16/18 instrument) was carried out in flowing air or 10% H2− N2 mixture at 298−1373 K with constant heating/cooling rate of 2 K/min. The total conductivity and Seebeck coefficient were measured simultaneously on the bar-shaped samples in flowing 10% H2−90% N2 mixture at 400−1173 K, decreasing the temperature by steps 50−80 K, and with equilibration at each step; the experimental setup is described elsewhere.21 The estimated experimental error in measured values did not exceed 5%. For selected compositions, oxygen partial pressure (p(O2)) dependence of the total conductivity was studied at 973−1173 K in isothermal conditions using H2−H2O−N2 mixtures. Oxygen partial pressure in atmosphere was monitored by yttria-stabilized zirconia (YSZ) solid-electrolyte sensor. Seebeck coefficient measurements were performed at typical temperature differences of 15−30 K; for all samples, the contribution of the offset voltage28 into the measured thermal voltage was found to be less than 3%. The thermal conductivity (κ = Dρcp) was calculated from the experimental results on thermal diffusivity (D) (Netzsch LFA 457 Microflash), specific heat capacity (cp) (Netzsch DSC 404 C), and density (ρ). As for total conductivity and Seebeck coefficient measurements, the experimental procedure included stepwise (50 K) change in temperature followed by dwell of 15−30 min for thermal equilibration of the sample. The estimated error in obtained

3. RESULTS AND DISCUSSION XRD analysis showed that all sintered Sr 1−xPrxTiO3±δ, Sr 1−1.5x Pr x TiO 3±δ (x = 0.05, 0.10, 0.20, 0.30), and Sr1−yTi0.8Nb0.2O3±δ (y = 0−0.10) samples are single-phase and have perovskite-type structure; typical XRD patterns are presented in Figure 1A. EDS inspection of the surface of as-prepared samples confirmed uniform distribution of the cations inside the grains and at the grain boundaries. The structure of Sr0.95Pr0.05TiO3±δ, Sr0.925Pr0.05TiO3±δ, and Sr1−yTi0.8Nb0.2O3±δ (y = 0−0.10) was identified as cubic, with space group Pm3̅m (Table 1), in agreement with literature data.24,29,30 The relative intensities of (100) and (210) reflexes compared to the main (100) peak differ for Nb- and Pr-containing materials, due to the different nature of the substituting cations, affecting the diffraction from atomic planes. Higher praseodymium concentration (x ≥ 0.10) for both Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ series results in lowering lattice symmetry to tetragonal, as indicated by the appearance of small addition peaks at 2Θ ∼ 38.2−38.4° and 50.9−51.1°, indexed as 211 and 213 reflections, and peak splitting at 2Θ ∼ 76.9−77.7° (Figure 1B,C). This structural transition is promoted by significant difference in ionic radii between host Sr2+ and Pr3+ cations causing the tilting of TiO6 octahedra.31 The refinement of XRD patterns for tetragonal perovskites was performed using I4/mcm space group (Table 1).31 Additional splitting of peaks at higher diffraction angles, as indicated by an arrow in Figure 1C, suggests further decrease in lattice symmetry for the sample with highest Pr content (Sr0.70Pr0.30TiO3±δ), which was successfully refined using orthorhombic space group Imma (Table 1); this is consistent with similar results reported for Sr0.7La0.3TiO3±δ.32 Still, this peak splitting was not observed for the corresponding sample with nominal compensation by A-site deficiency (Sr0.55Pr0.30TiO3±δ), which shows higher symmetry. 4598



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ing mechanism, realized via formation of electronic defects and, possibly, extended defects, such as rock-salt SrO layers characteristic for Ruddlesden−Popper (RP) phases, which act as a “sink” for excess strontium.24,34,35 Site occupancy for Sr1−xPrxTiO3±δ (0.05 ≤ x ≤ 0.30) and Sr1−yTi0.8Nb0.2O3±δ (0 ≤ y ≤ 0.06) can be represented as 4+ 3+ Sr 1−x−γ Pr x3+ Ti 1−x−2δ Ti x−2δ O 3−δ P × (SrO) γRP and P +2γ P −2γ 4+ 3+ Ti2δ Nb0.2O3−δP × (SrO)RP Sr1−y−αTi0.8−2δ α , correspondP+2α P−2α ingly, where δp is the oxygen nonstoichiometry of the perovskite-type layers. In the case of strontium deficient Sr1−1.5xPrxTiO3±δ and Sr0.90Ti0.8Nb0.2O3±δ, one expects the occupancies to be Sr 1 − 1 . 5 x Pr x3 + Ti 24−+δ Ti 24δ+ O 3 − δ and 3+ Sr0.90xTi4+ 0.8−2δTi2δ Nb0.2O3−δ, corresponding to nominal composition. However, the charge compensation in Pr-containing materials could be even more complex if one considers mixed Pr4+/Pr3+ state for praseodymium cations, as was demonstrated by XPS studies of similar materials, prepared in strongly reducing conditions.36 The latter may result in certain underestimation of the total oxygen content (Figure 2). Nevertheless, the total oxygen content shows a strong tendency to decrease from nominally stoichiometric to A-site deficient compositions, indicating an increase of the oxygen nonstoichiometry in the perovskite-type layers. Variations in the A:B cation ratio have a marked effect on the microstructure of ceramics prepared under similar conditions, as evidenced by SEM (Figure 3) and values of relative density (Table 1). In contrast to the Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ (0.05≤x ≤ 0.10) samples, which have quite similar microstructures (an example is shown in Figure 3A), the effect is more pronounced for higher donor-substitution level, namely, in Pr-containing series with x > 0.10 (Figure 3B,C) and Sr1−yTi0.8Nb0.2O3±δ (Figure 3D−F). Introduction of even minor cation deficiency, e.g., in Sr0.97Ti0.8Nb0.2O3±δ, results in significant grain growth and enhanced densification of ceramics compared to SrTi0.8Nb0.2O3±δ (Table 1). This behavior is in agreement with that observed for (Sr0.7La0.3)1−xTiO3±δ,32 Sr1±yTiO3±δ37, and Ba1±yTiO3±δ38 systems. Regardless of the microscopic mechanism of cation nonstoichiometry effect on the sintering process, the resulting difference in microstructure is expected to have an additional impact on thermoelectric properties in addition to that induced by the changes in point defect chemistry. Because the porosity may have a significant additional effect on electrical and thermal conductivities, the experimental data for σ and κ were corrected assuming that spherical pores are homogeneously distributed in the material.39 The temperature dependence of the total conductivity and Seebeck coefficient for Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ ceramic samples is shown in Figure 4. The negative sign of α confirms that conductivity is n-type in all prepared samples. It is usually assumed that under reducing conditions the concentration of n-type charge carriers (n or [Ti3+]) is determined by donor substitution level and oxygen deficiency in perovskite lattice.25 Even neglecting the possible presence of Pr4+, one should also take into account the contribution of A-site vacancies to charge neutrality, expressed as

The structural data presented in Table 1 point out that the increase in A-site cation deficiency results in a decrease of the unit cell volume for Pr-containing series, in correlation with the literature data on (Sr0.7La0.3)1−xTiO3±δ perovskites.32 For cubic Sr1−yTi0.8Nb0.2O3±δ perovskites, the changes in lattice parameter with strontium deficiency are, apparently, a result of complex defect formation and charge compensation mechanism, described below. In particular, isolated cation vacancies are expected to induce lattice expansion,33 whereas oxygen vacancies have opposite effects, and various effects from defect clustering also should not be neglected.32 The results of TG experiments corroborate the assumption21 that A-site substitution of Sr2+ by Pr3+ may result in oxygen excess in Sr1−xPrxTiO3±δ samples even under strongly reducing conditions; the same conclusion follows for A-site stoichiometric SrTi0.8Nb0.2O3±δ perovskite. The oxygen content in asprepared materials was estimated from the TGA data on complete oxidation in air at 1273 K, assuming negligible fractions of trivalent titanium in air ([Ti3+]→0) and that praseodymium is present in only the 3+ state. Thus, the results shown in Figure 2, were calculated as 3±δ=

m M mox ox

− nSrMSr − nPrMPr − n TiM Ti) MO

(1)

Figure 2. Total oxygen content in as-prepared Sr1−xPrxTiO3±δ, Sr1−1.5xPrxTiO3±δ, and Sr1−yTi0.8Nb0.2O3±δ ceramics, estimated from the thermogravimetric data on oxidation in air (see text).

where, m and mox are the weights of as-prepared sample at room temperature and completely oxidized sample, respectively; Mox, MSr, MPr, MTi, and MO are the atomic weights of completely oxidized sample, strontium, praseodymium, titanium, and oxygen, correspondingly; and nSr, nPr, and nTi are the number of Sr, Pr, and Ti (unity) atoms per unit formula, respectively. The original TG curves, corrected for baseline, are presented in the Supporting Information. Even for the smallest fraction of praseodymium A-site stoichiometric Sr1−xPrxTiO3±δ perovskites possess oxygen excess, suggesting a complex charge-compensat-

n + 2[V″Sr ] ≈ [D•] + 2[V •• O]

(2)

where [D•] is either [Pr•Sr] or [Nb•Ti]. The effect of A-site cation nonstoichiometry on electrical transport properties is clearly 4599



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Figure 3. SEM micrographs of fractured as-prepared ceramics: Sr0.85Pr0.10TiO3±δ (A), Sr0.70Pr0.30TiO3±δ (B), Sr0.55Pr0.30TiO3±δ (C), SrTi0.8Nb0.2O3±δ (D), Sr0.97Ti0.8Nb0.2O3±δ (E), and Sr0.94Ti0.8Nb0.2O3±δ (F).

distinguishable for moderate (0.05 ≤ x ≤ 0.10) and high (0.20 ≤ x ≤ 0.30) substitution levels (Figure 4A). For the former, the electrical conductivity decreases with increasing temperature, indicating metallic-type conduction in these samples. Considering the limiting case of γ = x/2 for site occupancy 4+ 3+ RP Sr1−x−γPr3+ x Ti1−x−2δP+2γTix−2δP−2γO3−δP × (SrO)γ , the concentration of charge carriers per formula unit can be calculated from TGA data on oxidation. In particular, the oxygen nonstoichimetry, shown in Figure 2, is interrelated with the calculated charge carrier concentration as [Ti3+] = 2δ + nPr. The results, shown in Figure 5, indicate that changes in [Ti3+] upon

introducing A-site deficiency in samples with x = 0.05 and 0.10 are in agreement with observed conductivity variations. Still, the charge carrier concentration can be underestimated in this case because of the possible presence of Pr4+ cations.36 Heavily substituted Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ series (0.20 ≤ x ≤ 0.30) show completely distinct conductivity behavior, including change from semiconductor-type in stoichiometric to metallic-type temperature dependence (Figure 4A) and higher Seebeck coefficients for A-site deficient materials (Figure 4B). Similar changes were induced by small difference in strontium content in SrTi0.8Nb0.2O3±δ and Sr0.97Ti0.8Nb0.2O3±δ ceramics, prepared in similar conditions, 4600



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localization of electronic charge carriers and, consequently, transition to thermally activated conductivity regime. Decreasing strontium content in the sample is expected to reduce this impact, preventing the formation of extended defects and thus providing conditions for delocalization of the charge carriers and metallic-type conductivity. Additional benefits may be also expected from the smaller unit cell volume for A-site deficient materials (Table 1) and corresponding higher overlap of Ti 3d orbitals.40,41 The most profound increase in Seebeck coefficient upon introducing A-site deficiency is observed for Prcontaining series with x = 0.30, consistent with improving crystal symmetry from orthorhombic Sr0.7Pr0.3TiO3±δ to tetragonal Sr0.55Pr0.3TiO3±δ. The higher local symmetry of TiO6 octahedra is known to enhance the density of states (DOS) effective mass of the carriers, thus increasing the Seebeck coefficient,42 which also might be the case for the studied materials. Among the factors responsible for the observed changes in electrical properties induced by A-site deficiency, one should also consider possible effects from significant microstructural variations in ceramic samples, in particular, porosity and grain size. Contrary to other ceramics, Sr1−xPrxTiO3±δ (x = 0.20− 0.30) and SrTi0.8Nb0.2O3±δ showed relatively high porosity (Table 1). For these donor-substituted titanates, defect reequilibration involves sluggish migration in the ation sublattice and defect distribution is kinetically nearly frozen below 1473 K.25,43 During cooling of the ceramics after sintering, high open porosity may promote oxygen exchange with the gas atmosphere, resulting in different distribution of cationic defects in porous and relatively dense material of the same composition. This mechanism is further confirmed by oxygen partial pressure-dependent total conductivity for porous samples (Figure 7), assuming that open porosity facilitates defect re-equilibration and considering gradual changes in charge compensation by electronic species and cation vacancies:25,44

Figure 4. Temperature dependence of the total conductivity (A) and Seebeck coefficient (B) for Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ ceramic samples.

Figure 5. Charge carrier concentration in Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ (x = 0.05, 0.10) estimated from thermogravimetric data on oxidation (see text).

while further increase in A-site deficiency up to 10% had no visible effect on the electrical properties (Figure 6A,B) within the sensitivity range of the measuring equipment.

x Sr Sr + 0.5O2 + 2e′ → SrORP + V″Sr

(3)

Figure 6. Temperature dependence of the total conductivity (A) and Seebeck coefficient (B) of Sr1−yTi0.8Nb0.2O3±δ ceramics.

It is worth noting that both electrical conductivity and the Seebeck coefficient were improved by nominal A-site deficiency, whereas one can expect a decrease in α when increasing the charge carrier concentration. Most likely, the observed trends should be attributed to a combined result of complex defect chemistry and differences in structural/ microstructural properties of the studied materials. High level of donor substitution leads to an increase in the number of insulating extended (planar) defects, which may induce

Figure 7. Total conductivity of Sr1−xPrxTiO3±δ, Sr1−1.5xPrxTiO3±δ, and Sr1−yTi0.8Nb0.2O3±δ ceramics versus oxygen partial pressure at 973− 1173 K. 4601



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For the dense A-site deficient samples, the conductivity shows plateaulike behavior (Figure 7), probably because response to redox changes is very sluggish, being determined by the frozen concentration of cation vacancies. Still, open porosity may reduce the time scale for redox processes by orders of magnitude as expected by extending the surface area available for surface exchange and/or by lowering the spatial dimension for diffusion controlled redox changes (LD) from the sample thickness in the millimeter range to individual grain sizes in the micrometer range. Note that the time scale for diffusion controlled processes varies as t ∝ LD2/D, D being the diffusion coefficient; this implies that changes from macroscopic sample thickness to grain size range may account for up to 6 orders of magnitude in time scale for diffusion controlled processes. In fact, Figure 7 shows significant redox changes only for quite porous samples, mainly for samples without A-site deficiency (Sr0.8Pr0.2TiO3±δ and SrTi0.8Nb0.2O3±δ). Thus, the role of A-site deficiency may be, at least partially, ascribed to its contribution to densification, with indirect consequences on the sluggish response to changes in the atmosphere. Comparison of high-temperature conductivity data for Sr1−xPrxTiO3±δ, Sr1−1.5xPrxTiO3±δ (x = 0.20−0.30), and Sr1−yTi0.8Nb0.2O3±δ (y = 0−0.03) suggests, however, that contribution from microstructural effects into the conduction mechanism may be different for Pr- and Nb-containing titanates. In particular, because for the samples sintered at 1773 K neither significant changes in microstructure nor additional reduction of Ti4+ are expected at 1100−1170 K, quite similar conductivity values for Sr 0.80 Pr0.20 TiO 3±δ, Sr0.70Pr0.30TiO3±δ, Sr0.55Pr0.30TiO3±δ, and even higher conductivity of Sr0.70Pr0.20TiO3±δ at high temperatures (Figure 4A) indicate that changes in crystal structure and defect chemistry, introduced by Pr3+ substitution and A-site nonstoichiometry, presumably are responsible for the observed trends in conductivity variations, as was previously suggested for (Sr,Pr)TiO3.21 Still, at lower temperatures (T < 873 K), one should not exclude possible impact of microstructural features such as high-ohmic depletion layers composed of acceptor-type states located at the grain boundaries.45,46 On the contrary, in the same conditions and in the whole studied temperature range, the electrical conductivity of Sr1−yTi0.8Nb0.2O3±δ (0.03≤y ≤ 0.10) is noticeably higher than that of SrTi0.8Nb0.2O3±δ (Figure 6A), while crystal structure remains cubic, indicating that microstructural difference between these ceramics may be a major reason for the variations in conduction mechanism. The similarity of transport properties measured for samples Sr0.97Ti0.8Nb0.2O3±δ, Sr0.94Ti0.8Nb0.2O3±δ, and Sr0.90Ti0.8Nb0.2O3±δ (Figure 6) is consistent with similar microstructures as well (Figure 3 E,F). The temperature dependence of the power factor for the studied materials is presented in Figure 8, illustrating the combined effect of electrical conductivity and Seebeck coefficient on thermoelectric performance. The largest power factor (PF) in Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ series was found for Sr0.90Pr0.10TiO3±δ by combining high electrical conductivity and relatively good Seebeck coefficient; this corresponds to one of the highest PF found for bulk polycrystalline oxide thermoelectrics (Figure 8A). Comparable power factors were found for A-site deficient Sr1−yTi0.8Nb0.2O3±δ (0.03 ≤ y ≤ 0.10) (Figure 8B). As already evident from σ/α data (Figures 4 and 5), the largest effect of Asite cation nonstoichiomentry on power factor is attained in the case of heavily substituted Sr1−xPrxTiO3±δ, Sr1.5−xPrxTiO3±δ (x

Figure 8. Temperature dependence of the power factor for Sr1−xPrxTiO3±δ, Sr1−1.5xPrxTiO3±δ (A), and Sr1−yTi0.8Nb0.2O3±δ (B) ceramics.

= 0.20−0.30), and Sr 1−y Ti 0.8 Nb 0.2 O 3±δ (y = 0−0.03). Introducing A-site deficiency in these materials is accompanied by a conduction mechanism change, represented by a transition from semiconductor to metallic-type behavior of the electrical conductivity with temperature, providing an increase in both electrical conductivity and Seebeck coefficient and consequently a large increase in power factor. Note that similar improvement of PF in the samples with metallic-type conduction was observed in the (Sr,La)TiO3 system, where high Seebeck coefficient was ascribed to the orbital degeneracy of Ti 3d-t2g conduction band and large effective mass of charge carriers.13 However, in the case of the materials studied in the present work, the impact of A-site cation deficiency on the power factor can be at least partially determined by the ceramics microstructure, as described above. Whatever the mechanism, at high temperatures the PF of Sr0.80Pr0.20TiO3±δ can be improved by A-site deficiency almost up to the level of that for Sr0.90Pr0.10TiO3±δ (Figure 8A), illustrating good prospects for tuning electronic properties of donor-substituted titanates by A-site nonstoichiometry level. Although both electrical conductivity and Seebeck coefficient of moderately substituted Pr-containing titanates (x = 0.05 and 0.10) were found to depend on A-site deficiency, thermal conductivity of these materials remains essentially unaffected 4602



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Taking into account the changes in oxygen nonstoichiometry, calculated from TGA results (Figure 2), a reasonable explanation for the lower thermal conductivity of A-site deficient materials in low and intermediate temperature ranges could be enhanced phonon scattering on oxygen vacancies. This assumption is also supported by previously observed significant reduction of thermal conductivity in oxygen-deficient lanthanum-substituted strontium titanate.15 Although formation of Ruddlesden−Popper phases was proven to be effective for suppressing thermal conductivity by improvement of phonon scattering at the SrO/(SrTiO3)n internal interfaces (e.g., ref 47), their possible impact on κ of A-site stoichiometric materials, studied in this work, appears to be negligible apparently because of low concentration. Still, one might attribute a small increase in lattice thermal conductivity from Sr0.97Ti0.8Nb0.2O3±δ to Sr0.90Ti0.8Nb0.2O3±δ to a decreased number of extended defects in the perovskite matrix (Figure 10A), although the related changes in κph are close to the experimental uncertainty. Other microstructural differences between A-site deficient and A-site stoichiometric compositions such as smaller grain size should result in lower thermal conductivity for A-site stoichiometric samples because of enhanced grain-boundary phonon scattering,48,49 which is contradictory to the obtained experimental results. Positive effect of the oxygen vacancies on lattice thermal conductivity in terms of thermoelectric performance is, however, hindered by the increased electronic contribution to the total thermal conductivity, as illustrated by Figure 11.

(Figure 9A). Lattice thermal conductivity (κph) was estimated from Wiedemann−Franz law as

Figure 9. Temperature dependence of the total (A) and lattice (B) thermal conductivity for moderately substituted Pr-containing ceramics.

κ ph = κ − LσT

(4)

where L is Lorenz number (2.44 × 10−8 V2 K−2) and shows similar trends for Sr1−1.5xPrxTiO3±δ with relatively low Pr content (Figure 9B). In contrast, A-site deficiency in heavily substituted Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ (0.20≤x ≤ 0.30) exerts important effects on thermal conductivity (Figure 10A), mainly on the lattice contribution (Figure 10B).

Figure 11. Lattice contribution to the total thermal conductivity.

As expected, for materials with metallic-type temperature dependence of the electrical conductivity, this contribution is higher at low temperatures. Still, heavily substituted A-site nonstoichiometric Sr0.55PrxTiO3±δ demonstrates total thermal conductivity at 420−620 K that is 1.2−1.5 times lower than that of nominal A-site stoichiometric Sr0.70Pr0.30TiO3±δ. At higher temperatures, the role of Umklapp scattering processes in the heat-transfer mechanism becomes dominant, while the influence of oxygen vacancies on thermal conductivity diminishes. Thus, A-site deficiency appears to be an effective means for suppressing lattice thermal conductivity of donorsubstituted strontium titanate at low and intermediate temperatures.

Figure 10. Temperature dependence of the total (A) and lattice (B) thermal conductivity for heavily substituted Pr- and Nb-containing series. 4603



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defect chemistry in the case of Pr-containing titanates, whereas for Sr1−yTi0.8Nb0.2O3±δ, perovskites microstructural evolution appears to be the most important factor. For both families of materials, introducing A-site deficiency represents a good strategy for improving thermoelectric performance at high donor substitution levels.

Summarizing the results discussed above, Figure 12 shows temperature dependence of dimensionless figure of merit ZT, calculated from obtained data on total conductivity, Seebeck coefficient, and thermal conductivity.

4. CONCLUSIONS Nominally stoichiometric and A-site cation deficient ceramic materials, Sr1−xPrxTiO3±δ, Sr1−1.5xPrxTiO3±δ (x = 0.05, 0.10, 0.20, 0.30) and Sr1−yTi0.8Nb0.2O3±δ (y = 0−0.10), were prepared by a conventional solid-state route and reduced at 1773 K in flowing 10% H2−90% N2. A-site cation nonstoichiometry was found to have a noticeable effect on the electrical properties, with less significant alteration of thermal conductivity. Large improvement in the power factor for heavily donor-substituted titanates was achieved because of favorable changes in the crystal structure, complex defect chemistry, and microstructural evolution. The relative role of each factor was found to be different in Pr- and Nb- containing titanates, and were also affected by temperature and substitution level. Presence of oxygen vacancies in A-site deficient materials, confirmed by thermogravimetry, provided a significant decrease in lattice thermal conductivity at low and intermediate temperatures. Increase in thermoelectric performance due to nominal A-site deficiency in titanates is ascribed mostly to improvement of power factor, whereas enhanced phonon scattering on oxygen vacancies was hindered by the growing contribution of the electronic part into the total thermal conductivity.

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ASSOCIATED CONTENT

S Supporting Information *

Additional thermogravimetry results for studied Pr- and Nbcontaining titanates. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author Figure 12. Temperature dependence of the dimensionless figure of merit ZT for Pr- (A) and Nb- containing (B) titanates.

*Department of Materials and Ceramic Engineering, CICECO, University of Aveiro, 3810-193 Aveiro, Portugal. Fax: +351234-370204. Tel: +351-234-370263. E-mail: [email protected].

For Sr1−xPrxTiO3±δ and Sr1−1.5xPrxTiO3±δ series, the ZT values increase on heating within the whole studied temperature range (Figure 12A), following the same trend as observed for many other A-site donor-substituted titanates.16,40,50,51 For A-site deficient Sr1−yTi0.8Nb0.2O3±δ (0.03 ≤ y ≤ 0.10), ZT reaches a maximum at ∼1100 K, where the changes in power factor and total thermal conductivity compensate each other (Figure 12B); similar behavior for donor-substituted titanates was also observed in ref 51. Although, among studied Prsubstituted titanates, maximal ZT corresponds to A-site stoichiometric Sr0.90Pr0.10TiO3±δ (Figure 12A), comparable thermoelectric performance (TEP) at high temperatures was achieved in heavily substituted Sr1−1.5xPrxTiO3±δ (x = 0.20 and 0.30), mainly due to large improvement in power factor with Asite deficiency. Sr1−yTi0.8Nb0.2O3±δ (0.03 ≤ y ≤ 0.10) also demonstrates superior TEP compared to cation-stoichiometric SrTi0.8Nb0.2O3±δ, sintered in the same conditions, e.g., approximately 1 order of magnitude higher at 1000 K. As discussed above, the reasons for ZT improvement can be at least partially related to the favorable changes in structure and

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the FCT, Portugal (projects SFRH/BPD/85619/2012, PEst-C/CTM/LA0011/2013, and FCT Investigator program, Grants IF/00302/2012 and IF/ 01072/2013). Financial support from the SNF-NCCR Manep and the DfG-SPP 1386 is greatly acknowledged. The authors are thankful to K. Galazka (EMPA) and S. Mikhalev, M.J. de Pinho Bastos (UA) for their experimental assistance.

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REFERENCES

(1) DiSalvo, F.G.. Thermoelectric Cooling and Power Generation. Science 1999, 285, 703−706. (2) Rowe, D. M. Thermoelectric Waste Heat Recovery as a Renewable Energy Source. Int. J. Innov. Energy Sys. 2006, 1, 13−23. (3) Tritt, T. M.; Subramanian, M. A. Thermoelectric Materials, Phenomena, and Applications: A Bird’s Eye View. MRS Bull. 2006, 31, 188−198.

4604



Article

structured EuTi1−xNbxO3−δ (x = 0.00; 0.02). Appl. Phys. Lett. 2012, 101, 033908. (24) Blennow, P.; Hagen, A.; Hansen, K. K.; Wallenberg, L. R.; Mogensen, M. Defect and Electrical Transport Properties of NbDoped SrTiO3. Solid State Ionics. 2008, 179, 2047−2058. (25) Moos, R.; Härdtl, K. H. Defect Chemistry of Donor-Doped and Undoped Strontium Titanate Ceramics Between 1000°C and 1400°C. J. Am. Ceram. Soc. 1997, 80, 2549−62. (26) Meyer, R.; Waser, R. Resistive Donor-Doped SrTiO3 Sensors: I, Basic Model for a Fast Sensor Response. Sens. Actuators, B 2004, 101, 335−345. (27) Rodríguez-Carvajal, J. Recent Advances in Magnetic Structure Determination by Neutron Powder Diffraction. Phys. B (Amsterdam, Neth.) 1993, 192, 55−69. (28) Iwanaga, S.; Toberer, E. S.; LaLonde, A.; Snyder, G. J. A High Temperature Apparatus for Measurement of the Seebeck Coefficient. Rev. Sci. Instrum. 2011, 82, 063905. (29) Durán, A.; Martínez, E.; Siqueiros, J. M. Structural Characterization and Ferroelectric Behaviour of Sr1−xPrxTiO3 Solid Solution at Room Temperature. Integr. Ferroelectr. 2005, 71, 115−120. (30) Irvine, J. T. S.; Slater, P. R.; Wright, P. A. Synthesis and Electrical Characterisation of the Perovskite Niobate-Titanates, Sr1−x/2Ti1−xNbxO3−δ. Ionics 1996, 2, 213−216. (31) Garg, R.; Senyshyn, A.; Boysen, H.; Ranjan, R. Structure of the Noncubic Phase in the Ferroelectric State of Pr-Substituted SrTiO3. Phys. Rev. B. 2009, 79, 144122. (32) Mori, M.; Nakamura, K.; Itoh, T. Crystal Structure and Anomalous Sintering Behaviour of (Sr0.7La0.3)1−xTiO3+δ Perovskites (0 ≤ x ≤ 0.12) Synthesized by the Pechini Method. J. Fuel Cell Sci. Technol. 2012, 9, 021007. (33) Ubic, R.; Tolman, K.; Khan, K.; Lundy, N.; Letourneau, S.; Kriven, W. M. Effective Size of Vacancies in Aliovalent Doped SrTiO3 Perovskites. J. Alloys Compd. 2013, 575, 239−245. (34) Balachandran, U.; Eror, N. G. Electrical Conductivity in Lanthanum-Doped Strontium Titanate. J. Electrochem. Soc. 1982, 129, 1021−1026. (35) Akhtar, M. J.; Akhtar, Z. U. N.; Jackson, R. A.; Catlow, C. R. A. Computer Simulation Studies of Strontium Titanate. J. Am. Ceram. Soc. 1995, 78, 421−428. (36) Yaremchenko, A. A.; Patrício, S. G.; Frade, J. R. Thermochemical Behavior and Transport Properties of Pr-Substituted SrTiO3 as Potential Solid Oxide Fuel Cell Anode. J. Power Sources. 2014, 245, 557−569. (37) Amaral, L.; Senos, A. M. R.; Vilarinho, P. M. Sintering Kinetic Studies in Nonstoichiometric Strontium Titanate Ceramics. Mater. Res. Bull. 2009, 44, 263−270. (38) Lin, M.-H.; Chou, J.-F.; Lu, H.-Y. The Rate-Determining Mechanism in the Sintering of Undoped Nonstoichiometric Barium Titanate. J. Eur. Ceram. Soc. 2000, 20, 517−526. (39) Zhang, L.; Tosho, T.; Okinaka, N.; Akiyama, T. Thermoelectric Properties of Solution combustion synthesized Al-doped ZnO. Mater. Trans. 2008, 49, 2868−2874. (40) Muta, H.; Ieda, A.; Kurosaki, K.; Yamanaka, S. Thermoelectric Properties of Lanthanum-Doped Europium Titanate. Mater. Trans. 2005, 46, 1466−1469. (41) Goodenough, J. B. Covalency Criterion for Localized vs Collective Electrons in Oxides with the Perovskite Structure. J. Appl. Phys. 1966, 37, 1415−1422. (42) Wang, Y.; Lee, K. H.; Hyuga, H.; Kita, H.; Ohta, H.; Koumoto, K. Enhancement of Thermoelectric Performance in Rare Earth-Doped Sr3Ti2O7 by Symmetry Restoration of TiO6 Octahedra. J. Electroceram. 2010, 24, 76−82. (43) Abrantes, J. C. C.; Labrincha, J. A.; Frade, J. R. Evaluation of SrTi1−yNbyO3+δ Materials for Gas Sensors. Sens. Actuators, B 1999, 56, 198−207. (44) Cumming, D. J.; Kharton, V. V.; Yaremchenko, A. A.; Kovalevsky, A. V.; Kilner, J. A. Electrical Properties and Dimensional Stability of Ce-Doped SrTiO 3−δ for Solid Oxide Fuel Cell Applications. J. Am. Ceram. Soc. 2011, 94, 2993−3000.

(4) Fujii, T.; Terasaki, I. Block-layer Concept for the layered cobalt oxide: A design for thermoelectric oxides. Proceedings of Chemistry, Physics, and Materials Science of Thermoelectric Materials: Beyond Bismuth Telluride, Traverse City, MI, August 17−21, 2002; Kanatzidis, M. G., Mahanti, S. D., Hogan, T. P., Eds.; Kluwer Academic/Plenum Publishers: New York, 2003; pp 71−87. (5) Terasaki, I.; Sasago, Y.; Uchinokura, K. Large Thermoelectric Power in NaCo2O4 Single Crystal. Phys. Rev. B 1997, 56, R12685− 12687. (6) Koumoto, K.; Terasaki, I.; Funahashi, R. Complex Oxide Materials for Potential Thermoelectric Applications. MRS Bull. 2006, 31, 206−210. (7) Terasaki, I.; Iwakawa, M.; Nakano, T.; Tsukuda, A.; Kobayashi, W. Novel Thermoelectric Properties of Complex Transition-Metal Oxides. Dalton Trans. 2010, 39, 1005−1011. (8) Weidenkaff, A.; Robert, R.; Aguirre, M. H.; Bocher, L.; Schlapbach, L. Nanostructured Thermoelectric Oxides with Low Thermal Conductivity. Phys. Status Solidi RRL 2007, 1, 247−249. (9) Hashimoto, H.; Kusunose, T.; Sekino, T. Influence of Ionic Sizes of Rare Earths on Thermoelectric Properties of Perovskite-type Rare Earth Cobalt Oxides RCoO3 (R = Pr, Nd, Tb, Dy). J. Alloys Compd. 2009, 484, 246−248. (10) Weidenkaff, A.; Robert, R.; Aguirre, M.; Bocher, L.; Lippert, T.; Canulescu, S. Development of Thermoelectric Oxides for Renewable Energy Conversion Technologies. Renewable Energy 2008, 33, 342− 347. (11) Maignan, A.; Hébert, S.; Pi, L.; Pelloquin, D.; Martin, C.; Michel, C.; Hervieu, M.; Raveau, B. Perovskite Manganites and Layered Cobaltites: Potential Materials for Thermoelectric Applications. Cryst. Eng. 2002, 5, 365−382. (12) Bocher, L.; Aguirre, M. H.; Logvinovich, D.; Shkabko, A.; Robert, R. R.; Trottmann, M.; Weidenkaff, A. CaMn1−xNbxO3 (x ≤ 0.08) Perovskite-Type Phases as Promising New High-Temperature ntype Thermoelectric Materials. Inorg. Chem. 2008, 47, 8077−8085. (13) Okuda, T.; Nakanishi, K.; Miyasaka, S.; Tokura, Y. Large Thermoelectric Response of Metallic Perovskites: Sr1−xLaxTiO3 (0 ≤ x ≤0.1). Phys. Rev. B 2001, 63, 113104. (14) Ohta, S.; Nomura, T.; Ohta, H.; Koumoto, K. HighTemperature Carrier Transport and Thermoelectric Properties of Heavily La- or Nb-Doped SrTiO3 Single Crystals. J. Appl. Phys. 2005, 97, 034106. (15) Yu, C.; Scullin, M. L.; Huijben, M.; Ramesh, R.; Majumdar, A. Thermal Conductivity Reduction in Oxygen-Deficient Strontium Titanates. Appl. Phys. Lett. 2008, 92, 191911. (16) Wang, H. C.; Wang, C. L.; Su, W. B.; Liu, J.; Zhao, Y.; Peng, H.; Zhang, J. L.; Shao, M. L.; Liu, J. C.; Yin, N.; et al. Enhancement of Thermoelectric Figure of Merit by Doping Dy in La0.1Sr0.9TiO3 Ceramic. Mater. Res. Bull. 2010, 45, 809−812. (17) Koumoto, K.; Wang, Y.; Zhang, R.; Kosuga, A.; Funahashi, R. Oxide Thermoelectric Materials: A Nanostructuring Approach. Annu. Rev. Mater. Res. 2010, 40, 363−94. (18) Muta, H.; Kurosaki, K.; Yamanaka, S. Thermoelectric Properties of Rare Earth Doped SrTiO3. J. Alloys Compd. 2003, 350, 292−295. (19) Liu, J.; Wang, C. L.; Su, W. B.; Wang, H. C.; Li, J. C.; Zhang, J. L.; Mei, L. M. Thermoelectric Properties of Sr1−xNdxTiO3 Ceramics. J. Alloys Compd. 2010, 492, L54−L56. (20) Shang, P.-P.; Zhang, B.-P.; Liu, Y.; Li, J.-F.; Zhu, H.-M. Preparation and Thermoelectric Properties of La-Doped SrTiO3 Ceramics. J. Electron. Mater. 2011, 40, 926−931. (21) Kovalevsky, A. V.; Yaremchenko, A. A.; Populoh, S.; Weidenkaff, A.; Frade, J. R. Enhancement of Thermoelectric Performance in Strontium Titanate By Praseodymium Substitution. J. Appl. Phys. 2013, 113, 053704. (22) Ohta, S.; Ohta, H.; Koumoto, K. Grain Size Dependence of Thermoelectric Performance of Nb-doped SrTiO3 Polycrystals. J. Ceram. Soc. Jpn. 2006, 114, 102−105. (23) Sagarna, L.; Shkabko, A.; Populoh, S.; Karvonen, L.; Weidenkaff, A. Electronic Structure and Thermoelectric Properties of Nano4605



Article

(45) Moos, R.; Härdtl, K. H. Electronic Transport Properties of Sr1−xLaxTiO3 Ceramics. J. Appl. Phys. 1996, 80, 393−400. (46) Karczewski, J.; Riegel, B.; Gazda, M.; Jasinski, P.; Kusz, B. Electrical and Structural Properties of Nb-Doped SrTiO3 Ceramics. J. Electroceram. 2010, 24, 326−330. (47) Lee, K. H.; Kim, S. W.; Ohta, H.; Koumoto, K. RuddlesdenPopper Phases as Thermoelectric Oxides: Nb-Doped SrO(SrTiO3)n (n = 1,2). J. Appl. Phys. 2006, 100, 063717. (48) Nolas, G. S.; Yang, J.; Goldsmid, H. J. Semiconductors and Thermoelectric Materials. In Thermal Conductivity: Theory, Properties and Applications; Tritt, T. M., Ed.; Kluwer Academic/Plenum Publishers: New York, 2004; pp 123−152. (49) Bocher, L.; Aguirre, M. H.; Robert, R.; Trottmann, M.; Logvinovich, D.; Hug, P.; Weidenkaff, A. Chimie Douce Synthesis and Thermochemical Characterization of Mesoporous Perovskite-Type Titanate Phase. Thermochim. Acta 2007, 457, 11−19. (50) Liu, J.; Wang, C. L.; Su, W. B.; Wang, H. C.; Li, J. C.; Zhang, J. L.; Mei, L. M. Thermoelectric Properties of Sr1−xNdxTiO3 Ceramics. J. Alloys Compd. 2010, 492, L54−L56. (51) Muta, H.; Ieda, A.; Kurosaki, K.; Yamanaka, S. Thermoelectric Properties of Lanthanum-Doped Europium Titanate. Mater. Trans. 2005, 46, 1466−1469.

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Effect of A-Site Cation Deficiency on the Thermoelectric Performance

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