Rare-Earth-Substituted Strontium Titanate: Insight into Local Oxygen

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Rare-Earth-Substituted Strontium Titanate: Insight into Local Oxygen-Rich Structures and Redox Kinetics Aleksey A. Yaremchenko,*,† Eugene N. Naumovich,‡ Sónia G. Patrício,† Oleg V. Merkulov,§ Mikhail V. Patrakeev,§ and Jorge R. Frade† †

CICECO−Aveiro Institute of Materials, Department of Materials and Ceramic Engineering, University of Aveiro, 3810-193 Aveiro, Portugal ‡ Fuel Cell Group, Thermal Processes Department, Institute of Power Engineering, Augustowka 36, 02-981 Warsaw, Poland § Institute of Solid State Chemistry, UB RAS, 91 Pervomayskaya Str., 620990 Yekaterinburg, Russia S Supporting Information *

ABSTRACT: Ln-substituted SrTiO3 is a promising material for energy conversion technologies such as thermoelectric generators and solid oxide fuel/electrolysis cells. In this study, formation of local structures enabling accommodation of excess oxygen in perovskite matrix of SrTiO3 and related redox behavior were assessed employing static lattice simulations in combination with experimental methods (XRD, SEM/EDS, XPS, TGA, and electrical measurements) using Sr0.90‑xLn0.10TiO3±δ (Ln = Ce, Pr; x = 0−0.10) as model systems. Although strontium-vacancy formation is found to be a preferable mechanism for donor compensation in oxidized Sr(Ln)TiO3, oxygen excess still can be accommodated by extended 2− defects quenched from high temperatures. Linear Ln3+ defect Sr ···Oi clusters and SrO shear planes characteristic of Ruddlesden−Popper phases are found to be the most probable extended defects enabling the accommodation of excess oxygen in oxidized titanates with Sr1−xLnxTiO3+δ cation stoichiometry. The presence of oxygen-rich local structures is shown to be strongly correlated with the faster redox kinetics and higher electrical conductivity critical for practical applications. Easy oxidation of reduced Sr1−xLnxTiO3±δ 2− (with electronic donor compensation) provide further evidence in favor of Ln3+ Sr ···Oi defect clusters as mechanism of excess oxygen accommodation.

1. INTRODUCTION Donor-substituted strontium titanate with perovskite-type structure exhibits outstanding phase and structural stability in a wide range of p(O2)−T conditions as well as suitable electrical properties and, therefore, is considered as a promising material for energy conversion at elevated temperatures. Potential applications include thermoelectric generators,1−3 interconnect materials for solid oxide fuel/electrolysis cells (SOFC/SOEC), 4,5 and ceramic components of SOEC cathodes6,7 and SOFC anodes.8−10 In the latter case, additional advantages of SrTiO3-based fuel electrodes include sulfur tolerance and resistance to carbon deposition8−10 which are important requirements for anodes of hydrocarbon-fueled SOFC. The magnitude of electrical conductivity (σ) is a critical parameter defining the thermoelectric efficiency of a material (thermoelectric figure of merit, ZT1) as well as its applicability in SOFC/SOEC electrodes and interconnects. Undoped SrTiO3 is a wide-gap semiconductor with comparatively low conductivity under both oxidizing and reducing conditions. Donor substitution in either sublattice (i.e., by Ln3+ in strontium sublattice or by Nb5+ in titanium sublattice) in © XXXX American Chemical Society

combination with thermal treatment under reducing conditions at very high temperatures results in an increase of n-type electronic conductivity by orders of magnitude. Depending on dopant content and pretreatment history, reported values of σ may vary from ∼1 to 300−400 S/cm at 700−900 °C (refs 8−10 and references therein). Often reported high conductivity values correspond, however, to a metastable state quenched from high temperatures and may be not redox-stable at operation temperatures (e.g., ref 11). Long-term stability may thus represent a major challenge for practical application. Although studied for nearly 4 decades, the defect chemistry of donor-doped strontium titanate is not established completely. Most commonly, two different mechanisms of donor compensation are considered for SrTiO3.10,12−15 Under oxidizing conditions, donor substitution is believed to be compensated by formation of vacancies in the strontium sublattice: Received: February 10, 2016

A

DOI: 10.1021/acs.inorgchem.6b00350 Inorg. Chem. XXXX, XXX, XXX−XXX

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Inorganic Chemistry (1 − 1.5x)SrO + x LnO1.5 + TiO2 →

Sr12−+1.5xLn 3x +Ti4 +O32 −

(1 − x)SrO + x LnO1.5 + TiO2 → Sr12−+1.5xLn 3x +Ti4 +O32 −*0.5x(SrO)RP

(1)

The formation of RP-like SrO planar faults (isolated or randomly distributed) in SrTiO3 with deviations from cation stoichiometry (Sr:Ti > 1.0) was demonstrated by highresolution electron microscopy (HREM)25−28 and supported by atomistic simulation calculations.28,29 An alternative mechanism of accommodation of excess oxygen was suggested for the Sr1−xLaxTiO3+δ system and involves the formation of crystallographic shear planes, i.e., oxygen-rich layers embedded parallel to a perovskite {110} planes.30−32 These layers have a structure similar to that occurring in La2Ti2O7 and have higher density of oxygen atoms compared to RP-like layers.30 At low levels of donor-type substitution, these planar defects presumably may occur as clusters or as short-range linear defects randomly distributed in the host perovskite, while the formation of compounds with ordered structures and general formula Srn‑4La4TinO3n+2 are characteristic for high substitution levels (4.5 ≤ n < 12); convincing evidence of the latter was provided by HREM and electron diffraction.30,31 At low donor-type substitutional level, the presence of extended defects, either RP-type or La2Ti2O7-like, cannot be detected by diffraction methods due to their low concentration and random distribution (Wadsley defects) in the perovskite matrix.21,30,31 Yet another mechanism of “self-compensation” is sometimes considered in the literature and may also explain the apparent oxygen excess in oxidized donor-substituted SrTiO3 with lanthanide additives.20,21 In this mechanism, Ln3+ substitutes in both strontium and titanium sublattices being simultaneously a donor- and acceptor-type dopant (depending on sublattice) and thus compensating itself; in this case oxygen content remains stoichiometric. Computer simulations20 predict that this mechanism may be favorable in the case of smaller rareearth cations (Ho, Yb, Lu), which possess intermediate ionic radii between the VI-coordinated B-site species and the XIIcoordinated A-site species. Evidence of partial redistributions of B-site species in A-site positions has also been reported for other perovskites such as A-site-deficient Ba1−x(Ce,Ln)O3‑δ33 or Sr1−x(Ce,Ln)O3‑δ.34 It was highlighted that the presence of oxygen-rich extended defects in the SrTiO3 structure increases the extent of Ti4+ → Ti3+ reduction under reducing conditions, in turn increasing the conductivity, and that oxygen access may play a critical role in both the structural and electrochemical properties.10 Understanding the true nature of extended defects is therefore extremely important for designing of modified strontium titanates with desired properties for energy conversion applications. The present work explores possible mechanisms of donor compensation and accommodation of excess oxygen in oxidized donor-substituted SrTiO3 employing static lattice simulations in combination with experimental techniques. Praseodymium- and cerium-substituted strontium titanates were selected as model materials. Both (Sr,Pr)TiO3±δ and (Sr,Ce)TiO3±δ systems were considered recently for possible applications as n-type thermoelectrics35−38 and as SOFC components.16,39,40 Nominal composition of the studied materials was Sr1‑x‑yLnyTiO3±δ, with concentration of donor dopant fixed at 10 at. % in strontium sublattice (y = 0.10), to study the comparatively diluted state of possible extended

Under reducing conditions, the donor is compensated via reduction of Ti4+ to Ti3+ (electronic compensation mechanism), with Ti3+ representing electrons (n-type charge carriers) contributing to the conduction process: (1 − x)SrO + x LnO1.5 + TiO2 −0.25xO2

⎯⎯⎯⎯⎯⎯⎯⎯⎯→ Sr12−+xLn 3x +Ti 3x +Ti14−+xO32 −

(2)

Assuming this, cation-stoichiometric compositions with high electrical conductivity should exist under reducing conditions, and insulating A-site-deficient compositions should exist under oxidizing atmospheres; a transition between two donor compensation mechanisms is expected to be accompanied by segregation or dissolution of secondary phases: Sr12−+xLn 3x +Ti14−+xTi 3x +O32 − + 0.25xO2 ox

XooY Sr12−+1.5xLn 3x +Ti4 +O32 − + 0.5xSrO red

(3)

red

Sr12−+1.5xLn 3x +Ti4 +O32 − XooY (1 − 0.5x)Sr12−+yLn 3y +Ti14−+yTi 3y +O32 − ox

+ 0.5x TiO2 + 0.25xO2

⎛ ⎞ x ⎜y = ⎟ ⎝ 1 − 0.5x ⎠

(5)

(4)

In practice, however, segregation of Ti-rich impurity on hightemperature reduction of single-phase A-site-deficient composition (eq 4) typically occurs only when strontium-vacancy concentration approaches a certain limit, i.e., ∼10% of A-sites (for instance, refs 12 and 16). This indicates that the reduced SrTiO3 lattice may tolerate coexistence of both A-site and oxygen vacancies up to certain thresholds. On the other hand, no evidence of SrO segregation in bulk samples of oxidized cation-stoichiometric compositions can be found in literature reports. This implies that these materials should be oxygen overstoichiometric under oxidizing conditions, and that there should be ways to accommodate excess oxygen in the crystal lattice. Accommodation of interstitial oxygen ions in closely packed SrTiO3 perovskite lattice (Frenkel disorder) or simultaneous formation of cation vacancies in both strontium and titanium sublattice to compensate for excess oxygen (Schottky disorder, similar to that occurring in LaMnO3+δ17) are typically considered unlikely.10,18−21 Instead, the formation of intergrowth structures (extended defects, or crystallographic shear planes) was suggested. There are two main possible mechanisms of accommodation of excess oxygen considered in the literature for the case of oxygen-excessive cationstoichiometric Sr1−xLnxTiO3+0.5x titanates. The first mechanism involves the compensation of donor doping by strontium vacancies, as in eq 3, accompanied however with the formation of SrO shear layers embedded parallel to {100} planes between the perovskite slabs of different thickness, similar to that occurring in layered Ruddlesden−Popper (RP) phases, Srn+1TinO3n+1 (n = 1, 2, 3, ...).13,21−24 Thus, SrO does not segregate as a separate phase, but both strontium and oxygen excess are accommodated within the perovskite matrix as extended defects: B

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SrTi1−zLnzO3±δ (z = 0.02 and 0.05) were prepared using the identical route with sintering in air at 1700 °C for 10 h. For preparation of porous Sr0.90‑xLn0.10TiO3±δ ceramic samples, precursor powders were further fired in air at 1300−1350 °C for 20− 30 h with repeated regrinding and final ball-milling. Disk-shaped samples were compacted uniaxially and sintered in air for 10 h at 1400 °C (cation-stoichiometric composition) or 1320 °C (A-site-deficient compositions). Figure S1 (see Supporting Information) illustrates microstructures of prepared porous samples. Sintered ceramics samples were polished and cut into rectangular bars (∼1.5 × 2.5 × 12 mm3; ∼0.22 g for dense samples and ∼0.16 g for porous samples) for electrical measurements. The density of prepared ceramics was calculated from the geometric dimensions and mass of polished samples. Measured density was ≥94% and 63−75% of the theoretical value in the case of dense and porous ceramics, respectively. Powdered samples for thermogravimetric studies were prepared by grinding dense sintered samples in a mortar. 2.3. Characterization. X-ray diffraction (XRD) patterns were recorded at room temperature using a Rigaku D/Max-B diffractometer (Cu Kα, 2Θ = 10−80°) using powdered samples. Microstructural characterization was performed by scanning electron microscopy (SEM, Hitachi SU-70) coupled with energy dispersive spectroscopy (EDS, Bruker Quantax 400 detector). Thermogravimetric analysis (TGA, Setaram SetSys 16/18 instrument, sensitivity 0.4 μg, initial weight of powdered samples ∼0.5 g) was carried out in flowing air or in a 10% H2−N2 mixture at 25−1100 °C with a constant heating/ cooling rate or isothermally as a function of time. The measurements of electrical conductivity as a function of oxygen partial pressure were performed by 4-probe dc technique in the p(O2) range from 0.21 down to ∼10−18 atm in O2−CO2 and CO−CO2 atmospheres using a specially designed YSZ chamber and bar-shaped ceramic samples with Pt wires as probes.45 Conductivity relaxation in isothermal air → 10% H2−N2 → air redox cycles was monitored using impedance spectroscopy (Agilent 4284A precision LCR meter) and bar-shaped samples with applied porous Pt electrodes. Oxygen partial pressure in gas mixtures was monitored by electrochemical yttria-stabilized zirconia (YSZ) oxygen sensors. Representative values of p(O2) in 10% H2−N2 mixture corresponded to ∼10−19 atm at 900 °C. 2.4. X-ray Photoelectron Spectroscopy. X-ray photoelectron spectroscopy (XPS) studies were performed at CEMUP (University of Porto, Portugal) using a Kratos AXIS Ultra HAS spectrometer with a monochromatic Al Kα radiation source (1486.7 eV) operating at 15 kV (90 W) in FAT (fixed analyzer transmission) mode. Measurements were made on a fractured sample surface with a takeoff angle of 90° at room temperature and total pressure ≤10−6 Pa. Dense ceramic samples were fractured immediately before insertion into the measurement chamber. Corrections for the sample charging were made using the C 1s band at 285 eV (from adventitious carbon) as an internal standard. XPS spectra were deconvoluted using XPSPEAK 4.1 software, employing a nonlinear least-squares fitting procedure with symmetric Gaussian−Lorentzian line shape and Shirley-type background subtraction. The intensity ratio between 3d5/2 and 3d3/2 spinsplit doublets was fixed to 3:2. Cation fractions were determined from the corresponding peak areas with a standard accuracy of ±10%. Pr6O11 (99.9%, Sigma-Aldrich) was analyzed as a reference material.

defects, with A-site deficiency x varying from 0 to 0.10. Note that the cases of Pr- and Ce-substitution are more complex when compared, for instance, to the Ln = La case. Both praseodymium and cerium may exist in mixed 3+/4+ oxidation state, with Ln3+ and Ln4+ possibly having preference for A- and B-sites, respectively, as it happens in Ce-substituted BaTiO3.41 Phase composition of ceramics materials was determined by X-ray diffraction and microstructural studies. The oxidation state of Ce and Pr cations was assessed using X-ray photoelectron spectroscopy. Attention is also given to the kinetics of redox changes on cycling between oxidizing and reducing atmospheres interrelated with defect chemistry mechanism; this was evaluated employing thermogravimetric analysis and electrical conductivity measurements.

2. METHODS 2.1. Static Lattice Simulations. Static lattice simulations were performed in GULP4 software using phenomenological potentials based on the Buckingham approach to interatomic interactions, and the core−shell model.42,43 Corresponding parameters are listed in Table S1 (see Supporting Information). The interatomic potentials of Pr4+ are not available for a given set of phenomenological potentials; therefore, simulations involving Ln4+ were limited to the Ce case. Lattice energies (Ecell) of simple oxides required for simulations were also calculated in GULP and are listed in Table S2. In the case of TiO2 (rutile), the value of Ecell was calculated as ° Ecell (TiO2 ) = Ecell (SrTiO3) − Ecell (SrO) − ΔGox (SrTiO3)

(6)

with the experimental value of ΔGox ° (SrTiO3) = −1.263 eV taken from the literature.44 The special procedure in this case was used due to a poor applicability of selected set of potentials for the rutile structure: the difference between experimental and simulated quantities exceeded 5%. The lattice simulations considered the formation of point defects (cation vacancies and interstitial oxygen) allowing preservation of crystal lattice electroneutrality in Ln3+-substituted SrTiO3, and possible cases of extended defects or structure reordering. All calculations were performed using supercells of appropriate dimensions, as described below. In order to assess the thermodynamic stability, the calculated lattice energies of supercells (SCs) were compared with the energies of the corresponding oxide mixtures: ΔEmix (SC) = Ecell (SC) − Emix (oxides) = Ecell (SC) −

∑ Ecell(oxides)

(7)

In all cases, the values of ΔEmix were normalized for a perovskite formula unit. 2.2. Synthesis and Ceramics Processing. Powders of Sr0.90‑xLn0.10TiO3±δ (Ln = Ce or Pr; x = 0.00, 0.05, 0.10) were prepared by conventional solid-state reaction route using SrCO3 (Sigma-Aldrich, ≥99.9%), TiO2 (Sigma-Aldrich, 99.8%), CeO2 (Sigma-Aldrich, 99.9%), and Pr6O11 (Sigma-Aldrich, 99.9%) as starting reagents. The precursor mixtures, weighed in appropriate proportions, were calcined consecutively at 900, 1100, 1200, and 1250 °C for 5 h at each temperature, with intermediate grinding, to attain high homogeneity. After subsequent ball-milling with ethanol, disk-shaped ceramic samples were compacted uniaxially and sintered in air at 1700 °C for 10 h. Hereafter, these samples are referred to as dense oxidized samples. Dense reduced ceramic samples were prepared by treatment of oxidized samples in a flow of 10% H2−N2 mixture at 1500 °C for 10 h. In the course of sintering/thermal treatment, Pt foil or a thick layer of the ceramic powder with the same cation composition was used as substrate in air and hydrogen-based atmospheres, respectively, to avoid possible reaction with alumina supports. In order to examine the possibility of incorporation of Ln3+ cations into the titanium sublattice, samples with nominal composition

3. RESULTS AND DISCUSSION 3.1. Phase Composition of Pr- and Ce-Substituted SrTiO3. Phase composition of as-prepared oxidized and reduced samples was determined employing XRD in combination with SEM/EDS. The results are summarized in Table 1. Selected XRD patterns are given in Figure S2 for oxidized ceramics and in Figure S3 for reduced materials. XRD pattern of oxidized cation-stoichiometric Sr0.90Ce0.10TiO3±δ revealed the presence of minor CeO2 impurity which dissolves on reduction. Phase impurities are absent not only in oxidized Sr0.85Ce0.10TiO3±δ, as expected for C

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significantly into the titanium sublattice of SrTiO3 under oxidizing conditions, at least in the synthesis and sintering routes employed in this work. On the contrary, the composition of secondary phases suggests that the substitution (if any) might occur in the strontium sublattice, and the A:B cation ratio in the perovskite structure was kept via segregation of SrCeO3 and Sr3Ti2O7 minority phases. 3.2. XPS: Oxidation States of Ce and Pr Cations in Oxidized Titanates. The oxidation states of cerium and praseodymium in the bulk of oxidized materials were determined by the analysis of XPS data collected from asfractured cross-sections of dense ceramic samples. Typical Ce 3d and Pr 3d spectra are presented in Figure 1; the corresponding core-level binding energies (BEs) and relative areas are summarized in Table S3.

Table 1. Phase Composition of As-Prepared Ln-Substituted SrTiO3±δ Ceramicsa phases nominal composition Sr0.90Ce0.10TiO3±δ Sr0.85Ce0.10TiO3±δ Sr0.80Ce0.10TiO3±δ Sr0.90Pr0.10TiO3±δ Sr0.85Pr0.10TiO3±δ Sr0.80Pr0.10TiO3±δ SrTi0.98Ce0.02O3 SrTi0.95Ce0.05O3 SrTi0.98Pr0.02O3 SrTi0.95Pr0.05O3

oxidized ceramics

reduced ceramics

C + CeO2 (1.3%) C C + TiO2 (1.3%) T C C + TiO2 (1.4%) C + CeO2 (0.9%) + SrCeO3 (0.6%) C + CeO2 (0.8%) + SrCeO3 (4.3%) C + Sr3Ti2O7 (1.3%) C + Sr3Ti2O7 (9%) + PrO1.833 (1.3%)

T T T + TiO2‑δ (n.a.) T T T + TiO2‑δ (n.a.)

a

C and T indicate the perovskite phase with cubic symmetry and with minor tetragonal distortion, respectively; numbers in parentheses are the ratios between the intensities of the strongest reflections of the impurity phase and perovskite phase.

cation vacancy charge compensation if one assumes prevailing trivalent lanthanide [Ln•Sr] ≈ 2[VSr ″ ], but also in reduced ceramics. Nominal composition Sr0.90‑xCe0.10TiO3±δ (x = 0.10) causes segregation of TiO2−δ, although the segregated fraction is low even in oxidized samples. Phase impurities were not detected by XRD in reduced Sr0.80Ce0.10TiO3±δ (Figure S3), but can clearly be seen and identified by microstructural analysis (Figure S4). Phase compositions of dense and porous ceramic samples fired in air were identical, except for porous Sr0.90Ce0.10TiO3±δ which contained a larger amount of the impurity phase (CeO2). The observations are in agreement with previously published reports on oxidized (Sr,Ce)TiO3±δ series.46,47 Phase analysis of Pr-substituted strontium titanates showed very similar results (Table 1 and Figures S2−S4) except that oxidized cation-stoichiometric Sr0.90Pr0.10TiO3±δ ceramics were single-phase as well. This difference between the (Sr,Ce)TiO3±δ and (Sr,Pr)TiO3±δ series can be explained by the tendency of cerium cations to retain higher oxidation state (4+) under oxidizing conditions. Tetravalent Ce4+ is less likely to be incorporated in A-site positions, possibly because its smaller ionic radius causes excessive deviations from the ideal tolerance XII XII factor (for comparison, rXII Sr2+ = 1.44, rCe3+ = 1.34, and rCe4+ = 1.14 Å).48 This results in an equilibrium Sr0.90Ce3+ 0.10−αTiO3±δ + αCe4+O2 phase mixture under oxidizing conditions, and is also consistent with dissolution of the minor CeO2 impurity in reduced Sr0.90Ce0.10TiO3±δ. On the contrary, praseodymium cations are already in mixed oxidation state in binary oxides in air, Pr6O11 (Pr3+:Pr4+ = 1:2) at room temperature, and Pr7O12 (Pr3+:Pr4+ = 1:0.75) at ∼800 °C, and tend to Pr2O3 at elevated temperatures (>1300 °C).49 Therefore, praseodymium can more easily adopt the 3+ oxidation state and substitute into the strontium sublattice in cation-stoichiometric composition. Attempts to prepare single-phase SrTi1−zLnzO3±δ materials failed. Ceramics with nominal composition SrTi1−zCezO3±δ contained CeO2 and SrCeO3 secondary phases, while XRD patterns of SrTi1−zPrzO3±δ samples showed the presence of Sr3Ti2O7 and PrOx impurities; the amount of impurity phases increased with z (Figure S2 and Table 1). These results indicate that cerium and praseodymium cations do not substitute

Figure 1. Deconvoluted high-resolution XPS spectra of Sr0.85Ln0.10TiO3±δ ceramics in (A) Ce 3d and (B) Pr 3d core-level regions, respectively. The inset illustrates a strong contribution of the oxygen Auger peak to the Pr spectrum. Open symbols correspond to the experimental data; overall simulated spectra are represented by gray lines. Common contributions of Pr3+/Pr4+ states are given by black lines, and the contributions assigned to 3+ and 4+ oxidation states are shown by blue and red lines, respectively.

The high-resolution Ce 3d XPS spectrum (Figure 1A) of oxidized Sr0.85Ce0.10TiO3±δ comprises two spin−orbit multiplets (v and u) corresponding to the 3d5/2 and 3d3/2 core levels, with the spin−orbit splitting of ∼18.3 eV. With the literature data taken into account,50−53 this photoemission spectrum can be interpreted as a superposition of Ce 3d spectra characteristic for Ce3+ and Ce4+ states. The main doublet at 885.2/903.6 eV (v′/u′) and v0/u0 satellite peaks tailing on the low-energy side were assigned to the 3d94f1 and 3d94f2 final states from the Ce3+ contribution, respectively. The doublet lines at 895.3/913.8 (v‴/u‴), 888.9/907.1 (v″/u″), and 882.9/901.2 (v/u) eV are characteristic for the Ce4+ contribution. In particular, the v‴/u‴ doublet corresponding to the 3d94f0 final state is typically D

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peak represents a fingerprint of the Pr4+ cations and is absent in the XPS spectrum of Pr2O3.16,56 The counterpart b″ peak in the 3d3/2 region could not be resolved due to overlapping with the strong oxygen Auger peak in the high-energy tail (see inset in Figure 1B). Besides these three doublets, the spectra exhibit one extra peak (t) at 957.7 eV assigned to the intra-atomic multiplet effect, as suggested by Ogasawara et al.56 The fractions of Pr4+ in (Sr, Pr)TiO3±δ were estimated using the method proposed by Borchert at al.:16,51

considered as a fingerprint of Ce4+ and is absent in spectra of Ce3+-based compounds.52−54 The v″/u″ and v/u doublets are due to a mixture of 3d94f1 and 3d94f2 configurations. Oxidized Sr0.90Ce0.10TiO3±δ exhibited a very similar XPS spectrum with slightly shifted binding energies and spin−orbit splitting of ∼18.5 eV (Table S3). Thus, the relative Ce4+ fraction (Table 2) was calculated from the peak areas using the following relationship:50−52 [Ce 4 +] [Ce 4 +] + [Ce3 +] area(v + v″ + v ″′ + u + u″ + u ″′) = total area

area(a″) [Pr 4 +] 1 = × 4+ c area(a′) [Pr ] + [Pr ] 3+

Here c is the constant calculated from the XPS data for the reference Pr6O11 sample with known fractions of Pr3+ and Pr4+. Under applied experimental conditions, c was equal to ∼0.19, and the relative contents of Pr4+ estimated from the XPS data are also listed in Table 2. Thus, XPS results confirm that cerium and praseodymium cations in oxidized (Sr, Ln)TiO3±δ ceramics have a strong preference for the 3+ oxidation state. The fraction of Ln4+ in the A sublattice is 11;31,32 instead, it was suggested that oxygen excess in compositions with n ≥ 12 may be accommodated in randomly distributed short-range linear defects within the perovskite framework.31,32 3.3.5. Continuous Defects in Perovskite Lattice. The formation of continuous lattice defects was simulated, aiming 2− to assess the possibilities of massive ordering of Pr3+ Sr ···Oi clusters. These studies were done employing 4 × 4 × 5 and 2 × 4 × 10 supercells. Different linear and planar cluster configurations (Figure 5) as well as formation of disordered 2− randomly distributed small 2Pr3+ clusters were considSr ···Oi ered. The results (Table 6) suggest that “infinite line” clusters should be more stable compared to other options, although the

cluster type 4 × 4 × 5 Supercell ”infinite line” “planar island” “infinite planar pattern” random distribution 2 × 4 × 10 Supercell ”infinite plane”

Ln

Ecell, eV

ΔEmix, eV

ΔEST mix, eV

Ce Pr Pr Pr Pr

−158.59 −158.60 −158.53 −158.53 −158.55

−1.07 −1.08 −1.00 −1.01 −1.02

0.06 0.06 0.14 0.13 0.12

Pr

−158.27

−1.04

0.10

ΔEmix and ΔEST mix are calculated with respect to equimolar (SrO + Ln2O3 + TiO2) and (SrTiO3 + Ln2O3 + TiO2) oxide mixtures, respectively. a

differences in formation energy values are very small. The stability of the “infinite line” configuration of oxygen interstitials was further evaluated with respect to nonideal occupancy of neighboring sites by Pr cations. These simulations were performed in a previously mentioned quasi-Monte-Carlo manner involving generation of two sets of lattices, with either 0 or 50% predefined placement of neighboring Pr cations along oxygen interstitials line and randomized distribution of H

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Taking into account the results of simulation and experimental data, one may suggest the following scenario: (1) Formation of oxygen-rich extended defects indeed occurs at very high temperatures close to sintering temperature. (2) These defects are probably not equilibrium at lower temperatures. Instead, Sr vacancy formation is likely to become a predominant defect chemistry mechanism. (3) This results in a segregation of SrO precipitates on the surface of ceramics during the thermal treatment at 1200−1400 °C, whereas precipitation of Sr-rich phase impurity in the bulk of ceramics is spatially arrested. (4) At temperatures ≤1000 °C, oxygen excess in cation-stoichiometric donor-substituted SrTiO3 is still accommodated by extended defects in a crystal lattice with kinetically frozen cation sublattice (i.e., a state quenched from high temperatures). For better comparison of different types of oxygen-rich extended defects or defect clusters which may form to compensate donor dopant in A-site of SrTiO3 lattice, the energies of Pr3+ incorporation were calculated in the same manner as ΔEST mix(SrTiO3 + Ln2O3 + TiO2) and ΔEmix(SrO + Ln2O3 + TiO2), but normalized per one praseodymium cation (Table 7). Although ordered structures with La2Ti2O7-like

remaining Pr cation in supercell. Energy optimization for 100 random supercells yielded 70% of stable configurations. The most stable configurations for each number of Pr3+ cations in the Oi line neighborhood were selected for analysis. The results of simulation are illustrated in Figure 6. Expectedly, the lattice

Figure 6. Energy of formation (ΔEmix) of Sr0.90Pr0.10TiO3.05 structure with “infinite line” defect with different fractions of praseodymium cations in the linear cluster (see text). Dotted line marks the energy of formation of A-site-deficient lattice with RP-like planar defect.

Table 7. Energy of the Pr3+ Incorporation for Sr0.90Pr0.10TiO3.05 stability increases with increasing fraction of praseodymium cations in the linear defect cluster. At the same time, one can note that the “infinite line” configuration offers lattice stability similar to that of the structure with RP-like planar defects when only 50% of Pr cations are actually ordered. 3.3.6. Comparison of Simulation Results: Most Favorable Donor Compensation Mechanisms. A comparison of the simulation results (Tables 3−6) indicates that donor compensation via formation of strontium vacancies accompanied by either mixed substitution (not supported by XRD) or with segregation of secondary phases (SrO, LnOx or Srn+1TinO3n+1) is the most favorable defect chemistry mechanism in the case of oxidized cation-stoichiometric Sr0.90Ln0.10TiO3+δ. Dissolution of rare-earth oxide in a highly stable SrTiO3 crystal lattice with formation of any type of extended defects or defect clusters that are able to accommodate excess oxygen seems to be unfavorable, at least at lower temperatures. One should point that firing of single-crystal and polycrystalline donor-doped cation-stoichiometric SrTiO3 at temperatures around 1200−1350 °C results in formation of SrO precipitates at the surface, and this is considered sometimes as a proof of strontium-vacancy compensation mechanism (e.g., refs 15, 65, and 66). The same phenomenon is characteristic for Pr- and Ce-substituted ceramics. For instance, thermal etching of polished dense oxidized Sr0.90Pr0.10TiO3+δ ceramics in air at 1400 °C results in formation of Ti-free (Sr,Pr)Ox precipitates on the surface (Figure S8). One should note, however, that, first, formation of Sr-rich precipitates occurs only on the surface of dense ceramics, but is not observed in the bulk. This is true for Sr0.90Pr0.10TiO3±δ and other oxidized Sr0.90Ln0.10TiO3±δ (Ln = La, Nd, Sm, Gd)36 materials prepared under similar conditions, although Sr0.90Ce0.10TiO3±δ is an exception as discussed above. Second, formation of SrO phase is not confirmed by SEM/EDS in the case of porous ceramics (probably due to limited resolution). Third, formation of such precipitates (Sr-rich after thermal treatment in oxidizing atmosphere and Ti-rich after etching in reducing atmosphere) is typical also for undoped SrTiO3 (e.g., refs 67 and 68).

donor-compensating defect

ΔEST mix per Pr cation, eV

ΔEmix per Pr cation, eV

Planar Defects RP-like structure 0.92 La2Ti2O7-like structurea 0.25 Defect Clusters in Cubic Perovskite “infinite line” 0.60 “planar island” 1.35 “infinite planar pattern” 1.31 random distribution 1.18 “infinite plane” 0.96 a

−10.44 −11.11 −10.76 −10.01 −10.05 −10.18 −10.04

Lattice stability is uncertain.

layers would be energetically more favorable, their formation in the case of low or moderate donor substitution results in a destabilization of simulated lattice (as discussed above) and also is ruled out by HR-TEM studies reported in the literature.31,32 2− It turns out that linear Pr3+ defect clusters and RP-like Sr ···Oi SrO shear planes are the most favorable extended defects enabling the accommodation of excess oxygen in oxidized donor-substituted strontium titanates with Sr1−xLnxTiO3+δ cation stoichiometry. One may also assume that RP-like structures may be in fact more favorable than estimated in the present work if they appear as more stable Sr3Ti2O7 intergrowths in a perovskite matrix (similar to that occurring in SrO−SrTiO3 system)25,28 rather than in isolated SrO shear planes as simulated here. 3.4. Redox Behavior. Whatever type of oxygen-rich defects may exist in oxidized donor-substituted SrTiO3-based ceramics, the reduction process is generally expected to occur in a similar manner. For cation-stoichiometric titanates with nominal Sr1−xLnxTiO3 composition and RP-like planar defects, the reduction process may evolve via the following reactions: (1) oxygen release from the perovskite matrix with formation of oxygen vacancies and trivalent titanium cations I

DOI: 10.1021/acs.inorgchem.6b00350 Inorg. Chem. XXXX, XXX, XXX−XXX

Article

Inorganic Chemistry

°C.69 Note that the corresponding range of spatial scale for diffusion-controlled processes (≈2(Dt)0.5) is about 0.2−1.2 μm within the time scale of 10 h; this indicates that redox kinetics may be affected by sluggish diffusion even at grain size scale. Neagu and Irvine71 proposed a mechanism of reduction of polycrystalline donor-substituted SrTiO3 ceramics which involves comparatively fast reduction of grain surface as a first step followed by a slow diffusion of oxygen vacancies from the surface into the bulk of grains until entire grains have been reduced. This mechanism is supported by the experimental data on electrical conductivity variations during redox cycling. Indeed, switching from an oxidizing to a reducing atmosphere under isothermal conditions results in a sharp initial increase of electrical conductivity of Sr0.90‑xLn0.10TiO3±δ samples (Figure 7)

Sr12−+1.5xLn 3x +Ti4 +O32 −*0.5x(SrO)RP −0.5δO2

⎯⎯⎯⎯⎯⎯⎯→ Sr12−+1.5xLn 3x +Ti14−+2δTi 24δ+O32 −− δ *0.5x(SrO)RP

(12)

(2) dissolution of RP-like defects with annihilation of Sr vacancies and formation of Ti3+ cations (a transition from cation vacancy compensation to electronic compensation of donor doping) Sr12−+1.5xLn 3x +Ti4 +O32 −*0.5x(SrO)RP → Sr12−+xLn 3x +Ti14−+xTi 3x +O32 −

(13)

(3) a combination of two processes Sr12−+1.5xLn 3x +Ti4 +O32 −*0.5x(SrO)RP −0.5δO2

⎯⎯⎯⎯⎯⎯⎯→ Sr12−+xLn 3x +Ti14−+x − 2δTi 3x ++ 2δO32 −− δ

(14)

For simplicity, these equations neglect the Ln4+ → Ln3+ transition and redissolution of CeO x in the case of Sr0.90Ce0.10TiO3±δ. The reduction process of oxygen-excessive titanates with La2Ti2O7-like planar defects can be expressed by similar equations. In the case of donor-doped SrTiO3 with defect clusters involving interstitial oxygen ions, the reduction can be imagined to occur similarly to reaction 12 but accompanied by dissociation of clusters and displacement of oxygen ions from interstitial to vacant regular sites. In the case of cation-deficient titanates with nominal composition Sr1−1.5xLnxTiO3, a reduction process also may occur via either oxygen release from the lattice with formation of oxygen vacancies and Ti3+ cations −0.5δO2

Sr12−+1.5xLn 3x +Ti4 +O32 − ⎯⎯⎯⎯⎯⎯⎯→ Sr12−+1.5xLn 3x +Ti14−+2δTi 32δ+O32 −− δ (15)

or through lattice reconstruction with segregation of titanium oxide impurity (eq 4), or as a combination of these two processes. Thus, in most scenarios, the overall reduction process should involve conjugate reconstruction of cation sublattice and reequilibration of oxygen sublattice with the gas phase. Strontium-vacancy diffusion was shown to be a predominant mechanism of cation sublattice re-equilibration in donorsubstituted SrTiO3.18,65 At the same time, cation diffusion in these materials was reported to be several orders of magnitude slower than oxygen diffusion.14,18,19 It is generally assumed that, at temperatures below 1200 °C, the mobility of strontium vacancies can be neglected, the cation sublattice is nearly frozen, and the oxygen vacancies are the only mobile species.13,14,69 This indicates that complete re-equilibration of the donor-substituted SrTiO3 lattice (i.e., eqs 4,13,14) in a reasonable time span is only possible at very high temperatures close to sintering temperatures.13,14 At temperatures below 1000 °C, the reduction process is essentially limited to reequilibration of oxygen sublattice (i.e., reactions 12 and 15). At the same time, re-equilibration in oxygen sublattice on reduction at