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Boosting thermoelectric performance by controlled defect chemistry engineering in Ta-substituted strontium titanate Aleksey A Yaremchenko, Sascha Populoh, Sonia Patricio, Javier Macías, Philipp Thiel, Duncan P. Fagg, Anke Weidenkaff, Jorge R. Frade, and Andrei V. Kovalevsky Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.5b01389 • Publication Date (Web): 29 Jun 2015 Downloaded from http://pubs.acs.org on July 4, 2015

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Chemistry of Materials

Boosting thermoelectric performance by controlled defect chemistry engineering in Tasubstituted strontium titanate

Aleksey A. Yaremchenko a, Sascha Populoh b, Sónia G. Patrício a, Javier Macías a, Philipp Thiel b, Duncan P. Fagg c, Anke Weidenkaff d, Jorge R. Frade a, Andrei V. Kovalevsky* a

a

CICECO – Aveiro Institute of Materials, Department of Materials and Ceramic Engineering, University of Aveiro, 3810-193 Aveiro, Portugal

b

Empa, Materials for Energy Conversion, Ueberlandstr. 129, CH-8600 Duebendorf, Switzerland

c

TEMA-NRD, Mechanical Engineering Department, Aveiro Institute of Nanotechnology (AIN), University of Aveiro, 3810-193 Aveiro, Portugal

d

Materials Chemistry, Institute for Materials Science, University of Stuttgart, Heisenbergstr. 3, DE-70569 Stuttgart, Germany

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

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Abstract

Inspired by recent research results that have demonstrated appealing thermoelectric performance of A-site cation-deficient titanates, this work focuses on detailed analysis of the changes in performance promoted by altering the defect chemistry mechanisms. The series of cation-stoichiometric SrTi1-xTaxO3±δ and A-site deficient Sr1-x/2Ti1-xTaxO3-δ compositions (0.05≤x≤ 0.30) with cubic perovskite-like structure were selected to demonstrate the defect chemistry engineering approaches, which result in promising electric and thermal properties. High power factors were observed in compositions where appropriate concentration of the charge carriers and their mobility were attained by presence of strontium- and oxygen vacancies and suppressed formation of the oxygen-rich layers. Noticeable deviations from stoichiometric oxygen content were found to decrease the lattice thermal conductivity, suggesting good phonon scattering ability for oxygen vacancies, vacant A-sites and oxygen-excessive defects, while the effect from donor substitution on the thermal transport was less pronounced. The obtained guidelines for the defect chemistry engineering in donor-substituted strontium titanates open new possibilities for boosting the thermoelectric performance, especially if followed by complementary microstructural design to further promote electrical and thermal transport.

Keywords: strontium titanate, thermoelectric performance, oxygen nonstoichiometry, thermal conductivity, Seebeck coefficient

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1. Introduction “Green” energy sources are rapidly becoming more and more important to meet the increasing energy demand, limited availability of the fossil fuels and various environmental issues, including CO2 emissions, global warming and impacts of other emissions on human health. One of the promising solutions is the thermoelectric conversion of waste heat into electricity by reliable, sustainable and scalable devices.1-4 Possible applications include waste heat recovery from power plants, industrial processes, furnaces, incinerators, gas heaters, engine exhaust streams, geothermal sources etc.5 Recent developments have pointed out good prospects of thermoelectric technology for power generation from concentrated solar irradiation6,7 and in the systems involving salinity gradient solar ponds, evacuated tube heat pipe solar collectors and biomass powered stoves.3 Many prospective applications, however, require the thermoelectric materials with high thermal and chemical stability, absence of toxicity and high natural abundance of the constituent elements. The efficiency of thermoelectric generation is limited by the Carnot efficiency and is governed by the figure of merit (ZT) of a candidate material, ZT = σ × α 2 × T / κ , determined by the electrical conductivity (σ), thermopower (Seebeck) coefficient (α) and thermal conductivity (κ). The numerator σ×α2 defines the power factor (PF), which primarily relates to electronic properties. In principal, the demand of performing thermoelectric materials for low- and intermediate temperature applications to a significant extent is covered by traditional Bi2Te3, Bi2Se3, PbTe –based thermoelectrics, intermetallic Zintl phases and half-Heusler alloys, skutterudites and some silicon-based alloys,8,9 with ZT values close to or higher than unity, allowing economically-feasible conversion efficiencies. Some of these materials can operate even at higher temperature, but lack stability in oxidizing atmospheres and are mostly toxic. Oxide materials have been considered as promising candidates for high-temperature thermoelectric applications, provided by their thermal and chemical stability, high natural abundance and less toxic composition.5,10 It should be noticed that high-temperature operation also yields higher Carnot efficiency of the conversion. A unique feature of the oxide materials is their redox flexibility, which allows controlled



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defect chemistry engineering by interaction with the atmosphere, in addition to heterovalent substitution effects. Yet, ZT values, obtained for the best oxide thermoelectrics are fairly below those demanded by most potential applications. Among oxides, donor-substituted strontium titanate shows great promises as n-type thermoelectric material due to it’s specific electronic structure, which can be tuned by introduction of structural defects, and prevailing lattice contribution to the thermal conductivity, enabling enhancements in phonon scattering by substitution and/or micro/nanoengineering approaches.5,11-18 Partial A-site substitution by rare-earth elements and B-site substitution by some transition metal cations, having stable oxidation state higher than four, are usually used as donor additives in SrTiO3 to attain reasonable electrical conductivity. In many cases, the substitution also suppresses the thermal conductivity by impurity scattering. Very recent works have highlighted good prospects for boosting thermoelectric performance in SrTiO3 –based materials by introducing an A-site deficiency,19,20 resulting in noticeably improved power factor and appearance of unusual “glass-like” behavior of the thermal conductivity. While pointing out high importance of the oxygen vacancies for both electrical and thermal transport, the work19 also discussed significant microstructural changes, promoted by A-site deficiency, which complicate the analysis of individual contributions from various defects. The work20 attributed low thermal conductivity of A-site deficient La-substituted strontium titanate exclusively to the presence of cation vacancies. At the same time, the observed reduction of the thermal conductivity in oxygen-deficient strontium titanates marks out the oxygen vacancies as efficient phonon scatterers.21 The latter is also supported by appealing performance of Sr0.9Ln0.1TiO3±δ materials, prepared in highly reducing conditions, where high ZT values were ensured by low thermal conductivity, apparently originating from the oxygen vacancies in the perovskite-type matrix.13 Based on these studies, understanding the exact role of the various defect types in the thermoelectric performance of strontium titanate-based materials still remains a challenge. In order to compare the impacts from defect structure, in the present work tantalum-substituted strontium titanate was selected as a model system, based on the following reasons. Firstly, deliberate B-site 4 Environment ACS Paragon Plus

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substitution should provide more insight regarding the effects caused by A-site deficiency, when strontium is the sole A-site cation. Secondly, tantalum is expected to have a clearer impact on the thermal conductivity compared to niobium, which is commonly-used B-site substituting cation for SrTiO3, due to the massdifference scattering. In particular, the literature data report contradictory effects from Nb substitution on the thermal conductivity of Sr(Ti,Nb)O3;22 the latter may complicate the interpretations for the influence of the defect chemistry features. The compositional design for shifting the prevailing defect types included two sets of materials, namely, cation-stoichiometric SrTi1-xTaxO3±δ and A-site deficient Sr1-x/2Ti1-xTaxO3-δ compositions (x=0.05-0.30). In this work we present a detailed analysis of the changes in thermoelectric performance, promoted by altering the defect chemistry mechanisms, focusing on relevant contributions of the electronic defects, substituting cations, cation and oxygen vacancies, and oxygen-rich planar defects.

2. Experimental Section. Series of stoichiometric SrTi1-xTaxO3±δ and A-site deficient Sr1-x/2Ti1-xTaxO3±δ compositions (x=0.050.30) were synthesized via a conventional solid state route, using SrCO3 (Sigma Aldrich, ≥ 99.9%), TiO2 (Sigma Aldrich, 99.8%) and Ta2O5 (Alfa Aesar, 99%) precursor powders. The basic processing routes are described in more details elsewhere.13,19,23 For the highest density of the samples, which is essential for the correct comparison of the thermoelectric properties of various chemical compositions, the ceramics were sintered by two-steps approach. The main densification step was conducted in air at 1973 K for 10 h. Presintered ceramics were further reduced in 10%H2-90%N2 atmosphere at 1773 K for 10 h. Subsequently, the prepared ceramic samples were ground into the fine powders for X-Ray diffraction (XRD), thermogravimetric (TG) and differential scanning calorimetry (DSC) studies. The total conductivity and Seebeck coefficient were measured on freshly-cut ceramic rectangular bars (~2×3×12 mm3). The sintered disc-shaped ceramics, polished down to ~1.00 mm thickness for removing possible surface contamination and providing uniform geometry, were used for the thermal diffusivity studies. For further assessment by SEM/EDS, the ceramic 5 Environment ACS Paragon Plus


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samples were polished and thermally etched. The experimental densities (ρ) of disk-shape ceramics were calculated from geometrical measurements and weighing. X-Ray photoelectron spectroscopy (XPS) was performed on the as-fractured sample surfaces. The room-temperature XRD patterns were recorded using a PANalytical X´Pert Pro diffractometer (Cu Kα). They were used for determination of the phase composition of prepared samples and calculation of the unit cell parameters by profile matching method in Fullprof software.24 Additional assessment of the phase purity and presence of any compositional inhomogeneities, which may be crucial for the thermoelectric performance, was performed by complementary SEM (Hitachi SU-70 instrument) and EDS (Bruker Quantax 400 detector) studies on fractured and polished ceramics. The total oxygen content of the studied materials was evaluated by TG; a detailed description of the corresponding procedure can be found elsewhere.19 The temperature dependence of the relative weight changes was measured in flowing 10%H2-N2 mixture at 2981373 K. The oxygen stoichiometry in as-prepared materials was estimated on complete oxidation in air at 1273 K, assuming 2+, 4+ and 5+ oxidation states for strontium, titanium and tantalum, correspondingly. X-Ray photoelectron spectroscopy (XPS) was performed at CEMUP (Porto, Portugal) in a Kratos AXIS Ultra HSA spectrometer equipped with monochromatic Al Kα radiation (1486.7 eV), operating at 150 W with a pass energy of 40 eV for regions ROI and 80 eV for survey. Corrections for sample charging effect were made by the reference of the adventitious carbon peak at 285 eV. Further experimental conditions were published previously.25 For Ti 2p, Ta 4f and Ta 4d, each spin-orbit doublet was fitted taking into account Scofield’s photoionization cross-section values.26 Atomic contents were determined with a standard accuracy of ±10% from the corresponding peak areas and normalized by the sensitivity factors provided by the manufacturer. To avoid possible oxidation and maintain the same conditions for all studied compositions, the ceramic samples were broken immediately prior to being placed to the measurement chamber; further XPS studies were performed on those fresh fractures. The total conductivity and Seebeck coefficient were measured at 300-1273 K in flowing 10%H290%N2 mixture. The measurement procedure included a stepwise decrease of temperature by 50-80 K, with up 6 Environment ACS Paragon Plus

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to 1 hour equilibration at each temperature. The maximal estimated error in the measured values corresponded to 3-5% for σ and 5-7% for α. Similar conditions were used for the determination of the thermal conductivity. It was calculated based on the experimental results on thermal diffusivity (D) (Netzsch LFA 457 Microflash), specific heat capacity (cp) (Netzsch DSC 404 C) and density as κ=Dρcp, with the estimated error less than 10%.

3. Results and discussion As the present study intents to elucidate the impact of structural defects features on relevant electrical and thermal properties, the question of phase purity and microstructural similarities of the prepared materials is an important issue. The results of room-temperature XRD studies of tantalum-substituted strontium titanate

110

(Fig. 1) indicate high solubility of Ta for both nominally A-site stoichiometric and Sr-deficient samples.

STT5 S85TT30

30

40

50

2Θ, °

60

221 300 310

210

100

20

TiO2 (rutile)

220

111

211

200

Intensity (a.u.)

A

70

0.396

B a, nm

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


SrTi1-xTaxO3±δ Sr1-x/2Ti1-xTaxO3±δ

0.394

0.392

0.390

0

0.1

0.2

x


0.3

80


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Fig. 1. Results of room-temperature X-Ray diffraction studies of the SrTi1-xTaxO3±δ and Sr1-x/2Ti1-xTaxO3-δ ceramics after sintering in reducing conditions: representative XRD patterns (A) and calculated lattice parameters (B).

Table 1 lists the chemical and phase composition of the prepared materials, together with corresponding abbreviations.

Table 1. Abbreviations, phase composition and density of SrTi1-xTaxO3±δ and Sr1-0.5xTi1-xTaxO3-δ ceramics.

Chemical

Abbreviation

Phase composition

composition

Density,

Relative

g/cm3

density*, %

SrTi0.95Ta0.05O3±δ

STT5

cubic perovskite

4.95

93.9


STT10

cubic perovskite

5.32

98.0


STT20

cubic perovskite

5.29

92.2


STT30

cubic perovskite

5.47

90.4

Sr0.975Ti0.95Ta0.05O3±δ

S975TT5

cubic perovskite

5.03

96.5

Sr0.95Ti0.90Ta0.10O3±δ

S95TT10

cubic perovskite

5.18

97.7

Sr0.90Ti0.80Ta0.20O3±δ

S90TT20

cubic perovskite

5.42

98.5

Sr0.85Ti0.70Ta0.30O3±δ

S85TT30

cubic perovskite +

5.33

93.1

minor TiO2 * -was calculated assuming that δ = 0

Separation of a minor secondary rutile phase was observed only for the S85TT30 sample, as a compensation for a significant deviation in A/B site ratio from the stoichiometric value (Fig. 1A). Interestingly, segregation involves the main B-site species (Ti), rather than the substituting species (Ta), suggesting that the actual substitution of Ta in B-site positions of the perovskite phase may at least up to 30%. Indeed, the ionic radii of 8 Environment ACS Paragon Plus

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Ta5+ (0.064 nm) and Ti4+ (0.0605nm) are similar,27 implying that the tolerance factor of Sr(Ti,Ta)O3±δ remains close to unity, providing that the SrTiO3-based perovskite unit cell is flexible to sustain other significant stresses, imposed by charge compensation or structural tuning. This allows phase-pure A-site nonstoichiometric materials up to at least 20% mol. of titanium substitution and Sr-deficiency, which corresponds to the nominal charge compensation for the donor Ta5+ additive. The compositional dependence of the lattice parameter closely follows a linear trend (Fig. 1B), in accordance with Vegard´s law, indicating the formation of solid solutions and a random distribution of tantalum in B-sites, except for the non-single phase S85TT30 sample. Combined SEM/EDS studies of the polished and fractured samples of all single-phase compositions did not reveal any compositional inhomogeneities, including those reported previously.28 The samples demonstrate fairly similar microstructures, with the grain size moderately increasing on substitution with tantalum. Typical SEM micrographs are shown in Fig. 2 (A-C). One should notice that the discussed variations in the grain size correspond to 3-15 µm. At this scale any noticeable effects from the grain size on the electrical and thermal conductivity are unlikely to contribute, being expected only for larger difference in average grain dimensions between the samples (at least, 1-2 orders of magnitude) or for submicron- and nanosized grain structures. This allows one to ascribe changes in thermoelectric properties to structural differences and structural defects, induced by cationic partial substitution and A-site deficiency, without accounting for microstructural effects, which may complicate the interpretation of the obtained results.19 The only exception is S85TT30 samples, for which the EDS analysis confirmed separation of titanium-rich phase (Fig. 2D), in agreement with the XRD data (Fig. 1, Table 1).



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(A)

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(B)

Sr Ti Ta (C)

(D)

Fig. 2. Representative SEM micrographs of the sintered samples, polished and thermally etched: STT10 (A), STT20 (B), S90TT20 (C), and EDS mapping results for S85TT30 (D) sample.

General guidelines regarding the defect chemistry of SrTi1-xTaxO3±δ and Sr1-0.5xTi1-xTaxO3-δ under reducing conditions can be obtained from thermogravimetric analysis. The estimates of the total oxygen content in as-prepared materials (Fig. 3) clearly show that the stoichiometric SrTi1-xTaxO3±δ materials possess


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detectable oxygen excess, while A-site nonstoichiometric compositions Sr1-0.5xTi1-xTaxO3-δ are oxygendeficient.

3.10 STT5 STT10 STT20 STT30

3.05

3±δ

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3.00 S975TT5 S95TT10 S90TT20

2.95

2.90

400

600

800 1000 1200

T, K Fig. 3. Variations of the total oxygen content with temperature, calculated from TG data.

Corresponding defect chemistry reactions and site occupancies can be described using standard Kröger-Vink notation.29 Since the close packed perovskite structure does not allow truly interstitial oxygen ions as point defects, possible charge compensation for oxygen excess in titanates should include the formation of extended planar rock-salt SrO layers, characteristic for Ruddlesden-Popper phases30,31 SrORP or Sr2Ta2O7-like layers

′′ ), similar to that occurring in Srn(Nb,Ti)nO3n+2 homologous series.32,33 accommodating excessive oxygen ( Olayer

( x 2)Ta 2O5 + SrO + (1 − x)TiO2 → (1 − β x / 2) SrSr× + (1 − x)TiTi× + ′′ + 3OO× + xTaTi• + ( β x / 2)VSr′′ + ( β x / 2) SrO RP + ((1 − β ) x / 2)Olayer

(1)

The contribution of the latter series is denoted as β and is expected to increase for higher tantalum content. These series can be viewed as a result of cutting the cubic perovskite structure along (110) direction followed by an insertion of additional oxygen.32



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In the cited works, those defects are described for the materials, prepared in oxidizing conditions. However, the obtained TG results strongly suggest that for A-site stoichiometric compositions their formation still takes place even in highly reducing conditions, providing the total oxygen content higher than three per perovskite unit cell. In terms of the crystal structure (ABO3+δ) it should be understood as SrORP and SrnTanO3n+2 planes alternating with the Sr(Ta,Ti)O3-δ layers of various thickness. Importantly, in the perovskite-type layers the presence of oxygen deficiency should not be excluded at the same time, although in terms of the overall oxygen content it is compensated by the oxygen-rich planes.13,19 Therefore, the nominal composition of the A-site stoichiometric compositions is represented as SrTi1-xTaxO3±δ, to account for different types of the oxygen-related defects. On the contrary, the defect chemistry of A-site nonstoichiometric materials is simpler and is expected to include the formation of electronic defects and oxygen vacancies due to reduction, while the excessive charge of Ta5+ is at least partially compensated by the strontium deficiency.

( x 2)Ta 2O5 + (1 − x 2) SrO + (1 − x)TiO2 → (1 − x / 2) SrSr× + ( x / 2)VSr′′ + + xTaTi• + (2δ )TiTi′ + (1 − x − 2δ )TiTi× + (3 − δ )OO× + δVO•• + (δ / 2)O2 ( g )

(2)

For these compositions the formulation Sr1-0.5xTi1-xTaxO3-δ in terms of the total oxygen content is relevant, except possibly for the nominal composition Sr0.85Ti0.7Ta0.3O3-δ, with Ti-rich precipitates. In this case, segregation is expected to yield even higher Ta contents, decreased A-site deficiency, and the corresponding changes in oxygen content in the resulting perovskite phase, as follows: (1 + α ) Sr0.85Ti0.7Ta 0.3O3−δ → αTiO2 + Sr0.85(1+α )Ti0.7−0.3α .Ta0.3(1+α ) O3±δ

(3)

The results of X-ray photoelectron spectroscopy (XPS), performed for selected compositions, provide further insight into the electronic-structure features of the prepared materials. Typical deconvoluted Ti 2p and Ta 4f and 4d high-resolution spectra are presented in Fig. 4. Corresponding core-level binding energies of the bands are summarized in Table 2. Peak fitting for the Ti 2p core level region (Fig 4A) shows that the XPS spectra consist of two 2p3/2 and 2p1/2 doublets. The main 2p3/2 and 2p1/2 doublets appear at 458.6 and 464.3 eV, respectively, with a spin-orbit splitting energy of 5.7 eV.


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Fig. 4. Deconvoluted high-resolution XPS spectra of Sr(Ti,Ta)O3±δ ceramics in (A) Ti 2p (B) Ta 4f and (C) Ta

4d core-level regions, respectively. Experimental points and the overall simulated spectra are represented by open symbols and gray lines, respectively. Spectra of Ti 2p were normalized by the Ti 2p3/2 peak intensity for easier comparison.

These binding energies and peak separation values are typical to those reported for TiO2 oxide,34 thus assigned to Ti4+ states. On the low energy-side, the additional 2p3/2 and 2p1/2 doublet with BEs peaks located at 456.3 and 462.3 eV respectively, match well with reported values for Ti3+ cations.34-36 It is also noteworthy that, for both composition sets, the binding energies of the main 2p3/2 sharp peak are slightly shifted to higher energies with increasing Ta content, likely due to the local structural distortions.



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Table 2 Core-level binding energies and relative areas of different components obtained by curve fitting of XPS spectra of Sr(Ti,Ta)O3±δ materials. Ti 2p Composition

STT5

STT10

STT20

S95TT10

Ta 4f

Ta 4d

BE, eVa

Area, %b

BE, eVa

Area, %b

BE, eVa

Area, %b

456.2 (0.80) 458.4 (1.06) 462.4 (0.58) 464.1 (1.97)

0.6 65.4 0.3 33.7

25.7 (1.00) 27.5 (1.05)

56 44

229.6 (3.70) 241.3 (4.12)

59 41

456.3 (1.15) 458.5 (1.20) 462.3 (1.04) 464.3 (2.11)

1.7 64.3 0.9 33.1

25.7 (1.17) 27.6 (1.25)

56 44

229.7 (3.89) 241.3 (4.30)

59 41

456.5 (1.35) 458.6 (1.16) 462.6 (1.03) 464.3 (2.10)

1.7 64.3 0.9 33.1

25.8 (1.11) 27.7 (1.18)

56 44

229.8 (3.94) 241.4 (4.25)

59 41

456.3 (1.11) 458.4 (1.20) 462.2 (1.11) 464.2 (2.06)

1.8 64.2 0.9 33.1

25.7 (1.19) 27.5 (1.25)

56 44

229.6 (3.87) 241.2 (4.10)

59 41

56 44

229.8 (4.01) 241.5 (4.24)

59 41

456.6 (0.95) 1.9 25.9 (1.34) 458.6 (1.44) 64.1 27.8 (1.47) 462.5 (1.05) 1.0 464.3 (2.30) 33.0 a Values in brackets refer to the FWHM (full-width at half maximum) of bands; b Area of each component relative to the total core-level peak area. S90TT20

In the case of tantalum, the high-resolution spectra were recorded in both 4f and 4d core level regions (Fig 4B and 4C, respectively). The Ta 4f spectrum can be resolved into one 4f7/2 and 4f5/2 doublet. The 4f5/2 peak at 27.7eV with a splitting energy of 1.9 eV from the 4f5/2 level indicates the presence of Ta5+ cations.36,37 The broader Ta 4d line, more reliable for identifying different oxidation states, also shows two symmetric spin orbit peaks, which confirm the assignment of one single oxidation state. The relative Ti3+ surface content and [Ta]total/[Ti]total atomic ratios are given in Table 3. As expected, titanium cations possess predominantly 4+ oxidation state, while the amount of Ti3+ is significantly lower and tends to increase with Ta substitution.


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Table 3 Surface atomic compositiona of Sr(Ti,Ta)O3±δ ceramics, estimated by XPS.

a b

Composition

[Ti3+]/[Ti]total, % a

[Ta (4f)/Ti] b

[Ta (4d)/Ti] b

[Ta/Ti]nominal

STT5

0.99

0.026

0.029

0.053

STT10

2.58

0.082

0.081

0.11

STT20

2.63

0.29

0.31

0.25

S95TT10

2.72

0.061

0.061

0.11

S90TT20

2.90

0.15

0.15

0.25

Determined by the areas of the respective peaks in the high resolution XPS spectra. Atomic ratios determined from XPS data

The estimated Ta/Ti atomic ratios, calculated from Ti 2p and Ta 4f and 4d core levels are quite similar or slightly higher compared to [Ta (4d)/Ti] ratio, due to the different escape depths. Though both ratios are relatively close to the nominal stoichiometry, these results still suggest that the surface tends to become Ti-rich in A-site deficient samples, eventually leading to onset of Ti-rich precipitates for the highest contents of Ta (Figs. 1 and 2). The XPS results indicate also that possible effects of partial reduction of tantalum cation to 4+ oxidation state on the defect chemistry of substituted strontium titanate can be neglected. As expected, altering the defect chemistry mechanisms results in noticeable changes of the electrical properties. Significant impacts from both substitution level and A-site deficiency on the electrical conductivity are well-illustrated by Fig. 5A. The conductivity is boosted by the introduction of the electronic defects, being, however, limited by other factors. Most of the studied materials show typical metallic-type behaviour of the electrical conductivity with temperature, as observed for other donor-substituted strontium titanates. By analogy with other donor-additives, substitution of Ti4+ with Ta5+ promotes formation of Ti3+ and filling of the

n-type conduction band, also confirmed by a negative sign of the Seebeck coefficient in the whole studied temperature range (Fig. 5B).



STT5 S975TT5 STT10 S95TT10 STT20 S90TT20 STT30 S85TT30

A

2.5 2.0 1.5 1.0 0.5

-350

B

-300

α, µV×K-1

3.0

log σ (S×cm-1)

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-250 -200 -150 -100

0.0 6

9

12

15

18

21

24

-50

400

600

104/T, K-1

800

1000

1200

T, K

Fig. 5. Temperature dependence of the electrical properties of SrTi1-xTaxO3±δ and Sr1-0.5xTi1-xTaxO3±δ ceramics: total conductivity σ (A) and Seebeck coefficient α (B).

Another contribution to the charge carrier concentration is expected from the partial reduction and formation of the oxygen vacancies ( VO•• ) under strongly reducing conditions during sintering:

( x 2)Ta 2O5 + SrO + (1 − x)TiO2 → SrSr× + ( x 2 + 2δ )TiTi′ + xTaTi• + + (1 − 1.5 x − 2δ )TiTi× + (3 − δ )OO× + δVO•• + (0.25 x + δ / 2)O2 ( g )

(4)

This known mechanism implies that, in the simplest case, the concentration of n-type carriers (n) should roughly correspond to:

n = [ TiTi′ ] ≈ [ TaTi• ]+2[ VO•• ]

(5)

However, in both cation-stoichiometric SrTi1-xTaxO3±δ and A-site deficient Sr1-0.5xTi1-xTaxO3±δ series the conductivity significantly increases only up to 10% mol. of Ta, while showing a tendency of further decrease for higher tantalum content. A likely reason for that is the relatively high energy of the 5d-derived states of tantalum, resulting in much less effective d-p hybridization with oxygen ions38 if compared to titanium. Thus, Ta5+ acts rather as an impurity in the conducting planes, as expected for metallic conduction. In other words, titanium substitution by tantalum in SrTiO3 results in a trade-off impact of the increasing concentration of the 16 Environment ACS Paragon Plus

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charge carriers and their scattering at Ta5+ sites on the electronic transport. The highest electrical conductivity corresponds to 10-20% mol. Ta, also depending on the A-site deficiency. The heavily-substituted STT30 sample shows distinct semiconducting-type behaviour of the electrical conductivity with temperature, if compared to the other materials and, in particular, to the S85TT30 composition, having the same substitution level. Moreover, the values of the electrical conductivity for STT30 are significantly low. These observations indicate that additional structural features affecting the electronic transport are also present in this case. As shown above, A-site design approach allows fine tuning of the defect chemistry mechanism by shifting the prevailing defect types. The contribution of those defects into the relevant physical properties can be, therefore, determined. Thus, one may attribute a rapid decrease in the electrical conductivity in the sequence STT10>STT20>>STT30 to the prevailing impact from RP and oxygen-excessive defect planes, as suggested by TG data (Fig. 3). A significant amount of these defects is likely to induce the localization of the electronic charge carriers and, consequently, transition to the thermally-activated conduction regime for STT30. For this highest amount of tantalum, introduction of A-site deficiency, however, restores metallic-type conduction and high conductivity values, as observed for S85TT30. Additional conductivity boosting in S85TT30 sample should be provided by partially reduced TiO2 inclusions,39 where presence of Ti3+ cations enables fast electronic transport, in opposite to insulating oxygen-excessive planes in STT30. In general, oxygen deficiency in the perovskite layers is favorable for the electronic transport, as it provides higher concentration of the n-type carriers. Hence, all A-site deficient compositions demonstrate higher electrical conductivity than their stoichiometric counterparts, with the maximum observed for S95TT10 and S90TT20 samples. In addition, there are indications that A-site deficiency itself is likely to promote the reduction and generation of oxygen vacancies.40 In particular, when a large number of A-sites are vacant, the oxygen ion might not be required to pass through the bottleneck defined by the A-B-A triangle in the perovskite ABO3 structure, thus facilitating the oxygen diffusion, which is the rate-determining step in the reduction of bulk ceramics.40



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The concentration of n-type charge carriers should increase significantly upon substitution from STT5 to STT10 sample (Table 3), in agreement with the conductivity data. Using [Ti3+]/[Ti]total atomic ratio, measured by XPS, and total conductivity one can roughly estimate the carrier´s mobility, to obtain more guidelines for the impacts from defect chemistry. The results of XPS studies, performed at room temperature, are expected to be relevant at higher temperatures as well, as the defects distribution in donor-substituted titanates is kinetically nearly frozen below 1473 K.19,25,41 The concentration of the charge carriers may be expressed as:

n = [Ti 3+ ] /[Ti ]tot × (1 − x) × N fu / Vuc

(6)

where x is the substitution level, Nfu is the number of formula units per unit cell and Vuc is the unit cell volume. The mobility (µ) of the charge carrier is given by:

µ = σ / ne

(7)

where n is the carrier charge. The results of calculations are presented in Table 4.

Table 4 Charge carrier concentration, mobility and effective mass at 600 K, estimated from XPS/total conductivity/Seebeck coefficient data using Eqs. 6-8.

#

Composition

n, cm-3

µ, cm2×V-1×s-1

m*/m0#

STT5

1.57×1020

4.19

1.20

STT10

3.85×1020

5.34

1.48

STT20

3.46×1020

2.55

0.99

S95TT10

4.07×1020

7.23

1.50

S90TT20

3.82×1020

8.63

1.08

- m0 – free electron mass

The calculated n values give a more clear picture for the substitution level effect on the charge carrier concentration than [Ti3+]/[Ti]total ratios, derived from the XPS (Table 3). As the tantalum concentration 18 Environment ACS Paragon Plus

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increases, the number of titanium sites, available for the reduction, effectively decreases. Thus, above ~1015% mol. Ta content no further noticeable increase in the charge carrier concentration is observed. This explains why the highest electrical conductivity is observed for STT10, S95TT10 and S90TT20 samples, compared to other compositions. The estimated mobilities are comparable to those directly measured for lanthanum- and niobium-substituted strontium titanate.14 The mobility increases only slightly from STT5 to STT10, suggesting that for low and moderate substitution the electrical conductivity is mainly determined by increasing the number of the charge carriers upon substitution. Further addition of tantalum results rather in the formation of oxygen-excessive defects than in generation of mobile charge carriers. The formation of extended defects (Eq. 1) decreases the carrier´s mobility in the STT20 sample by promoting their scattering. A striking feature is the highest mobility, estimated for S95TT10 and S90TT20 compositions. Logically, this might be attributed to structural specifics of the A-site deficient compositions, which may facilitate the electronic transport. Additional studies, including modelling of the electronic band structure and direct measurements of the Hall mobility and charge carrier concentration, may be required for deeper assessment of this behaviour. Nevertheless, the obtained results clearly show that the mobility of the charge carriers is greatly enhanced by the introduction of the A-site deficiency, likely by decreasing the amount of scattering (SrO)RP –type and oxygen-excessive planes. The presence of structural defects, where excess oxygen is accommodated, apparently does not affect the Seebeck coefficient, as indicated by similar values of α for stoichiometric and A-site deficient series. This observation is consistent with the fact that the Seebeck coefficient is rather determined by the concentration of the charge carriers and the electronic band structure of the perovskite matrix. The absolute α values significantly decrease for both stoichiometric and A-site deficient materials up to 10% mol. of Ta, due to an increase in the charge carrier’s concentration (Table 4). Further changes in the Seebeck coefficient are rather moderate, provided by smaller variations in n. The S85TT30 sample shows a deviation from the main trends, likely due to presence of the second phase and corresponding shifts in the perovskite matrix composition. Another important parameter, responsible for the changes in Seebeck coefficient, is the charge carrier effective 19 Environment ACS Paragon Plus


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mass (m*). It can be roughly estimated by assuming a parabolic band model in the energy independent scattering approximation,15,42 as:

m* =

3αeh 2 8π 2 k B2T (π / 3n) 2 / 3

(8)

where e, h and kB are the carrier charge, Planck´s and Boltzmann constants, correspondingly. The highest Seebeck coefficient of STT5 is provided by low n, coupled with moderate carrier effective mass (Table 4). Stoichiometric and A-site deficient compositions show comparable m* values and similar Seebeck coefficients. One should notice that the apparently higher charge carrier mobility in A-site deficient S90TT20 compared to S95TT10, as discussed above, may be attributed to noticeably lower effective mass (Table 4). Indeed, in quasi-classical approach of the semiconductor theory the carrier mobility can be represented as:43

µ=

eτ m*

(9)

where τ is a phenomenological scattering time to account for the scattering of the electrons by impurities and phonons. However, this approach fails to explain the differences in mobility and effective mass, observed for stoichiometric compounds, most probably due to a more complex scattering mechanism, as can be already expected from their defect chemistry features. It is generally accepted that the thermoelectric power factor increases with the carrier effective mass, as the gain in squared Seebeck coefficient is normally larger than decrease in mobility.16 This is consistent with the results, presented in Fig. 6. Largest power factor was observed for S95TT10 and STT10 samples with highest estimated m* (Table 4), for the former reaching a maximum of ~1.42 mW×m-1×K-2 at 470 K, which is certainly among the best PF values observed in SrTiO3-based thermoelectrics. The improvements in PF due to A-site deficiency are higher at lower temperatures. A comparison of power factors of the stoichiometric and A-site deficient compositions at “intermediate” temperatures reveals significant gains of ~50% from STT5 to S975TT5, ~20% from STT10 to S95TT10 and even ~220% from STT20 to S90TT20 at 800 K. In the case of 30% mol. of tantalum the improvement amounts to ~88 times, provided by the change from semiconductor- to metallic type behavior of the electrical conductivity due to absence of scattering oxygen-excessive layers and 20 Environment ACS Paragon Plus

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presence of oxygen deficiency. These two oxygen-related defects are, thus, very important for engineering high-power-factor donor-substituted titanates.

1600

Power factor, µW×m−1×K-2

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


STT5 STT10 STT20 STT30

1400 1200

S975TT5 S95TT10 S90TT20 S85TT30

1000 800 600 400 200 40 20 0 400

600

800

1000

1200

T, K

Fig. 6. Temperature dependence of the power factor.

Moderately substituted samples in both stoichiometric and A-site deficient series show quite similar total thermal conductivities (Fig. 7A), in agreement with the results obtained previously for (Sr,Pr)TiO3±δ materials.19 The thermal conductivity is composed of the contributions from the lattice (κph) and electronic counterparts (κel). In many oxide materials, including titanates, the lattice thermal conductivity prevails over electronic, and the temperature dependence of κ at higher temperatures is roughly proportional to the inverse temperature. The lattice contribution can be estimated from Wiedemann-Franz´ law:

κ ph = κ − LTσ

(10)

where L is the Sommerfeld value (2.45·10-8 W×Ω×K-2) of the Lorenz number.44 The results are shown in Fig. 7B. 21 Environment ACS Paragon Plus



5

S975TT5 S95TT10 S90TT20 S85TT30

4 3

5

κph, W×m−1×K-1

6

κ, W×m−1×K-1

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

2

4 3 2

B

A 1

400

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600

800

1000

1200

1

400

T, K

600

800

1000

1200

T, K

Fig. 7. Temperature dependence of the total (A) and lattice thermal conductivity (B).

In simplified case of the Callaway model, the lattice contribution to the thermal conductivity can be presented as:45-47

κ ph

1 = 3

ωmax

∫ C (ω )v s

g

(ω ) 2τ (ω )dω

(11)

0

where ω is the phonon frequency, νg is phonon group velocity, Cs – spectral heat capacity and τ is the phonon relaxation time. According to Matthiessen's rule, the lattice thermal conductivity is determined by a combination of point defects scattering (τPD), grain boundary scattering (τGB) and Umklapp phonon-phonon interactions (τU): 1 / τ = 1 / τ PD + 1 / τ U + 1 / τ GB

(12)

Umklapp scattering is less affected by the defect chemistry and starts to dominate at high temperatures, leading to similar thermal conductivities of the materials having different composition and defect structure. Grain boundary scattering can be roughly accounted as:47 22 Environment ACS Paragon Plus

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τ GB =

d vg

(13)

where d is the average grain size and vg is the phonon group velocity. One expects a negligible contribution from this scattering mechanism in the case of the studied materials, as it only becomes noticeable in nanostructured ceramics.47 At lower temperatures the thermal conductivity decreases upon Ta substitution for both stoichiometric and A-site deficient series (Fig. 7B). If simplified to the case of alloying on a single crystallographic site, the relaxation time for point defects scattering can be presented as: −1 = τ PD

m r Vω 4 (∑ f i (1 − i ) 2 + ∑ f i (1 − i ) 2 ) 2 m r 4πv p v g i i

(14)

where V is the volume per atom, vp – phonon phase velocity, fi – the fraction of atoms with mass mi and radius ri, residing on a site with average mass and radius m and r , respectively.47,48 Tantalum is around 3.5 times heavier than titanium and is expected to contribute to the reduction in the thermal conductivity upon substitution due to a high mass contrast. Presence of oxygen and A-site vacancies will also significantly contribute to the deviations of the mass and size of the crystallographic sites from the average, with corresponding impact on the thermal transport. However, assuming 10% mol. Ta-containing materials and minor impact of the size mismatch term, the mass difference term f i (1 −

mi 2 ) due to the substitution can be m

estimated as roughly equal to 0.38, while the total contribution from A-site deficiency and oxygen nonstoichimetry in the case of S95TT10 composition does not exceed ~0.1. In other words, one should expect a larger reduction of the lattice thermal conductivity from STT5 to STT10, if compared to the difference between STT10 and S95TT10 samples, which is contradictory to the obtained results (Fig. 7B). However, similar

thermal

conductivities

were

also

observed

for

SPS-consolidated


and

SrTi0.90Ta0.10O3±δ titanates.28 This may indicate that a complex defect chemistry of the donor-substituted strontium titanate encompasses additional factors affecting the thermal transport, including size-dependent effects and local distortions of the crystal lattice due to presence of VO•• and VSr′′ defects. Another issue is that vacancies may scatter phonons in a different manner, compared to atoms.20,49 23 Environment ACS Paragon Plus


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In moderately substituted compositions the presence of Sr- and oxygen deficiency results in a noticeable decrease of the lattice thermal conductivity, likely due to the reasons described above. The highest

κph for STT5 and STT10 apparently correlates with the total oxygen content close to stoichiometric (Fig. 3), due to the lowest expected concentration of the lattice imperfections and extended defects. For heaviersubstituted compositions the κph(T) behaviour shows much weaker temperature dependence and significantly lower values of κph. Recent work on A-site deficient SrTiO3-based thermoelectrics20 attributed those changes to the appearance of low energy “rattling” vibrations modes that dissipate heat, promoted by vacant crystallographic sites. In present work, similar behaviour was also observed for A-site stoichiometric STT20 and STT30 samples (Fig. 7B). Although these compositions still may possess minor oxygen and/or A-site cation deficiency in the perovskite layers,13,19 it is not expected to be close to the one of S90TT20 and S85TT30. At the same time, the lattice thermal conductivities of STT20 and S90TT20 samples are very similar (Fig. 7B). These trends suggest that any reasonable deviation from the ideal oxygen stoichiometry (3) in Sry(Ti,Ta)O3±δ is favourable for suppressing the lattice thermal transport. For clarity, Fig. 8 illustrates this behaviour by a green area, which corresponds to the effects, imposed on the lattice thermal conductivity by oxygen-related defects and, most likely, A-site cation vacancies.


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5

κph, W×m−1×K-1

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


600 K

3.002 3.007

4 2.975

3

2.950 2.927 3.041

3.077

2 0.00

0.10

0.20

0.30

x (Ta) Fig. 8. Lattice thermal conductivity of SrTi1-xTaxO3±δ (blue closed circles) and Sr1-0.5xTi1-xTaxO3±δ (red open circles) ceramics, estimated at 600 K. The numbers indicate total oxygen content, obtained from TG data. The green area illustrates the κph reduction due to oxygen and cation vacancies.

It has to be noticed that within the concept of the present work, the separate contributions provided by oxygen and cation vacancies for κph reduction cannot be determined, since all studied oxygen-deficient compositions also possessed cation deficiency. One may expect to distinguish them by further reduction of A-site stoichiometric compositions. A noticeable decrease in the thermal conductivity due to oxygen vacancies was already confirmed for Sr1−xLaxTiO3−δ films,21 as an evidence for their contribution to the lattice thermal transport. The concentration of heavy substituting cation is another contributing factor (Eq. 14). It is believed that a simple atomic substitution could not result in appearance of glass-like lattice vibrations,49 which one may assume for x≥0.2, considering significant changes in κph magnitude and temperature dependence, and literature data.20 In a disordered crystal the glass-like contribution to the thermal conductivity may be linked to the presence of random, non-central distortions of the lattice,49 likely being promoted by the strains exerted by vacant sites. Thus, it is expected that significantly lower lattice thermal conductivity of x≥0.2 compositions is rather provided by the oxygen and strontium vacancies, together with oxygen excessive layers, than by donor 25 Environment ACS Paragon Plus


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substitution with tantalum. Titanium exsolution and corresponding enrichment of the S85TT30 main phase with strontium may be, thus, responsible for a higher κph of this composition compared to STT30, due to the “defects release” promoted by phase separation, namely the expected suppressing of A-site vacancies (Eq.3). One should notice that the moderate difference in the samples densification (Table 1) is not expected to affect the discussed correlations between the chemical composition, defect structure and thermal conductivity of the studied materials. As an example, assuming that spherical pores are homogeneously distributed in the material and applying Maxwell correction for porosity50 without accounting for any structural factors, an increase in porosity from STT10 to STT30 sample should result in only ~9% decrease in the thermal conductivity, while the experimentally observed difference (Fig. 7A) between these compositions corresponds to 45-60%. Thus, it should be noticed that the actual decrease in the thermal conductivity from STT10 to STT30 samples due to compositional and structural factors is slightly lower than shown in Figs. 7,8, taking into account this additional porosity contribution. The cumulative impact of the electrical and thermal transport properties on the thermoelectric performance is illustrated by the temperature dependence of the dimensionless figure of merit ZT (Fig. 9). Attractive ZT values, attained by defect chemistry engineering, were observed for A-site nonstoichiometric oxygen-deficient S95TT10 and S90TT20 samples and amounted to 0.30 and 0.28 at 1000 K, and to 0.37 and 0.35 ant 1230 K, respectively. The key impact of the defect chemistry, resulting in high thermoelectric performance, was, however, different for these compositions. In S95TT10 the major boost is promoted by the high power factor, likely provided by high carrier mobility and effective mass. In S90TT20 the main parameter, responsible for high ZT, is the thermal conductivity, efficiently suppressed by high lattice-defects concentration. This comparison clearly shows high potential of the defects engineering in donor-substituted strontium titanates for breaking the coupling between electrical and thermal properties towards enhanced thermoelectric performance.


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0.4

S975TT5 S95TT10 S90TT20 S85TT30

0.3

ZT

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.2 0.1 0.0 400

600

800

1000

1200

T, K Fig. 9. Thermoelectric figure-of-merit of SrTi1-xTaxO3±δ and Sr1-0.5xTi1-xTaxO3±δ ceramics.

The obtained ZT values are significantly higher than those previously reported for bulk polycrystalline Sr(Ti,Ta)O3 (0.17 at 752 K)28 and Sr(Ti,Nb)O3 (0.09 at 800 K,51 0.22 at 1050 K,52 0.165 at 900 K,17 0.23 at 1000 K19). Still, the thermoelectric performance of Sr(Ti,Nb)O3, consolidated by SPS (0.33 at 1050 K),5 hot pressing (0.35 at 1000 K),18 or deposited as epitaxial thin films (0.2953, 0.354 and 0.3755 at 1000 K) is comparable or higher than that observed in the present work for the best materials; corresponding data is also plotted in Fig. 1S (Supporting Information). Thus, defect structure engineering should be complemented with specific processing routes, allowing even more extensive defects creation, including favourable micro- and nanostructural features. It should be noticed that detailed HRTEM studies would be helpful to assess the formation and evolution of the extended defects in reduced titanates. At the same time, discussed defect chemistry mechanism is strongly supported by the relevant literature data on similar materials, including the results of high-resolution electron microscopy and computer simulations,56-60 and reasonably confirmed by the TG results. Thus, in the present work we focused more closely on the comparative analysis of the impacts from various expected defects on the thermoelectric properties, with particular emphasis on the interrelation between substitution level and A-site composition as a driving force for tuning the defect structure. 27 Environment ACS Paragon Plus


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4. Conclusions Cation-stoichiometric SrTi1-xTaxO3±δ and A-site deficient Sr1-x/2Ti1-xTaxO3±δ compositions (x=0.050.30) were prepared by conventional solid state route and treated under highly-reducing conditions to assess the impact of various structural defects on the thermoelectric performance. The compositional design allowed fine tuning of the defect chemistry by shifting the prevailing defect types from extended oxygen-rich defects to oxygen nostoichiometry on introducing the A-site deficiency in perovskite layers. The effect of the donorsubstitution level on electrical properties has a complex nature, governed by the interrelation between generation of the charge carries and their scattering at the substituted sites and defects, accommodating excessive oxygen. Presence of both cation- and oxygen vacancies were found favourable for higher mobility of the charge carriers. The Seebeck coefficient was mostly determined by the charge carrier concentration, being less affected by other defects. High power factors, reaching ~1.42 mW×m-1×K-2 at 470 K for Sr0.95Ti0.90Ta0.10O3±δ, were attributed to the low concentration of oxygen-excessive layers and presence of oxygen deficiency in perovskite matrix. “Glass-like” lattice thermal conductivity was observed for both A-site deficient and stoichiometric compositions at x≥0.20, and was attributed mostly to the presence of A-site and oxygen-related defects, while the apparent impact of the B-site substitution was less pronounced. The described strategies for the defects engineering resulted in attractive ZT values over a large temperature range, observed for A-site nonstoichiometric oxygen-deficient Sr0.95Ti0.90Ta0.10O3-δ and Sr0.90Ti0.80Ta0.20O3-δ, amounted to 0.30 and 0.28 at 1000 K, and to 0.37 and 0.35 ant 1230 K, respectively.

Supporting Information Comparison of the thermoelectric performance for selected materials, prepared in this work, with the literature data on the best B-site donor-substituted titanates. This information is available free of charge via the Internet at http://pubs.acs.org/. 28 Environment ACS Paragon Plus

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Acknowledgements This work was supported by the FCT Investigador program (grants IF/00302/2012 and IF/01072/2013), and was developed within the scope of projects BPD/75943/2011, SFRH/BD/91675/2012, and project CICECO-Aveiro Institute of Materials (ref. UID/CTM/50011/2013), financed by national funds through the FCT/MEC and when applicable co-financed by FEDER under the PT2020 Partnership Agreement. Financial support from the Competence Centre Energy and Mobility (CCEM), the Swiss Federal Office of Energy (BfE) within the HITTEC project is also greatly acknowledged. The authors are thankful to Prof. Carlos Sá (CEMUP) for performing XPS studies and helpful discussion of the results, M.J. de Pinho Bastos (UA) for her assistance with XRD analysis and to Dr. S.M. Mikhalev for his technical support.



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ToC

0.4

Sr1-x/2Ti1-xTaxO3±δ

1200 K

ZT

0.3 0.2 0.1

SrTi1-xTaxO3±δ

0.0 5

κph, W×m−1×K-1

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

600 K

3.002 3.007

4 2.975

3

3±δ 2.950

2.927 3.077 3.041

2 0.00

0.10

0.20

0.30

x (Ta)


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