Influence of the Oxygen Content on the Electronic Transport Properties

Article pubs.acs.org/JPCC

Influence of the Oxygen Content on the Electronic Transport Properties of SrxEu1−xTiO3‑δ Leyre Sagarna,† Sascha Populoh,† Andrey Shkabko,† James Eilertsen,† Alexandra E. Maegli,† Roland Hauert,‡ Matthias Schrade,§ Lassi Karvonen,† and Anke Weidenkaff*,† †

Solid State Chemistry and Catalysis and ‡Nanoscale Materials Science, Empa − Swiss Federal Laboratories for Materials Science and Technology, Ü berlandstrasse 129, CH-8600 Dübendorf, Switzerland § Centre for Materials Science and Nanotechnology, University of Oslo, Gaustadalléen 21, NO-0349 Oslo, Norway ABSTRACT: The electronic transport properties of perovskite-type oxides with promising thermoelectric properties can be tuned by cationic substitutions as well as by creating oxygen vacancies. Here, two series of SrxEu1−xTiO3‑δ (x = 0.00, 0.03, 0.25, 0.75, 0.97) samples with different oxygen stoichiometries are synthesized using various reductive atmospheres. Oxygen deficiencies in the series produced in highly reducing atmosphere are confirmed by thermogravimetric analysis. Electrical resistivity (ρ) and Seebeck coefficient (S) data reflect the strong influence of the oxygen nonstoichiometry on the transport properties. Samples from the strongly reduced series possess up to 5 orders of magnitude lower electrical resistivities than samples from the less reduced series, confirming that oxygen vacancies act as electron donors in the former. Their Seebeck coefficients strongly decrease with the increase of the charge carrier concentration (lower ρ). Pronounced S(T) peaks are observed, and are attributed to strong hybridization of the localized Eu 4f states with the delocalized Ti 3d states. The presence of optically active defect states is confirmed by UV−vis spectroscopy.

1. INTRODUCTION Europium titanate, EuTiO3, crystallizes in a cubic ABO3 perovskite structure with Pm3̅m symmetry at room temperature, where the Eu2+ ions occupy the A-site and the Ti4+ ions occupy the B-site. The strongly localized Eu 4f spins adopt a Gtype antiferromagnetic (AFM) ordering below TN = 5.3 K.1 Katsufuji and Takagi2 reported strong spin−lattice coupling in EuTiO3, and observed that the dielectric constant increases with decreasing temperature and decreases further at the AFM transition temperature. Akamatsu et al.3 proved that the AFM ordering happens via superexchange between the Eu 4f spins mediated through the 3d states of nonmagnetic Ti4+ ions. The possibility of controlling the dielectric constant by a magnetic field and the magnetization by an electric field stimulated a variety of studies on EuTiO3,4−6 and it became a highly interesting material in multiferroic research.7−9 Other europium perovskite compounds, like EuMoO3,10 have shown strong correlations between the electric conduction and a large magnetic moment. The thermoelectric properties of Lasubstituted EuTiO311 and the large Seebeck coefficient (S) of EuTiO3 (−1084 μV/K at 268 K)12 suggested the application of EuTiO3-related materials in the field of thermoelectricity. Partial anionic substitution was carried out by Sagarna et al.,13 and the crystal structure and thermoelectric properties of EuTi(O,N)3±δ were reported. The employment of low-priced elements is desired for energy related applications. Therefore, Eu should be avoided or used in lower proportions. Kato et al.14 investigated the thermoelectric properties of (Sr1−xEux)Ti0.8Nb0.2O3 (x = 0, 0.1, 0.2, 0.5, and 0.8). No significant dependence upon the Eu content was observed in the © 2014 American Chemical Society

electronic transport properties, but the lowest thermal conductivity was obtained for x = 0.5. The solid solution Sr1−xEuxTiO3 is appealing due to the similarities between EuTiO3 and SrTiO3, which might enlarge the field of application.15 The ionic radii of Eu2+ and Sr2+ are almost identical, and consequently, EuTiO3 and SrTiO3 have nearly the same lattice parameters, a = 3.904 Å and a = 3.905 Å, respectively.16,17 An important difference between both perovskites is that SrTiO3 shows a structural phase transition from cubic Pm3m ̅ to tetragonal I4/mcm at the transition temperature Ts = 105 K,18 whereas the analogous structural transition associated with a similar rotational instability of the oxygen octahedra was suggested recently for EuTiO3 at higher temperatures. The TS of EuTiO3 has been discussed by different authors and temperatures from 235 to 298 K were reported.19−21 However, Bessas et al.22 studied the lattice dynamics of EuTiO3 and found no phase transition. They suggested that the lattice instabilities close to room temperature occur due to the delocalization of Eu atoms from their positions. The phase diagram of Sr1−xEuxTiO3 was determined by Guguchia et al.15 as a function of x in order to explore the large difference between the structural transition temperatures. They found a nonlinear dependence of TS with x. Other studies on the solid solution Sr1−xEuxTiO3 have focused on the crystal structure, microstructure and dielectric properties.23 In addition, the oxidation states of the cations were studied. For Received: January 15, 2014 Revised: March 18, 2014 Published: March 21, 2014 7821

dx.doi.org/10.1021/jp500480c | J. Phys. Chem. C 2014, 118, 7821−7831

The Journal of Physical Chemistry C

Article

and weight of the pellets, resulting in an average value of 65% ± 2% of the corresponding theoretical densities. A second series of SrxEu1−xTiO3‑δ samples was produced in situ during the high-temperature measurement of the electrical resistivity (ρ) and Seebeck coefficient (S), in a dynamic (50 mL/min) 5% H2/95% Ar atmosphere. This series exhibited higher oxygen stoichiometry than the 30HAr samples. In order to differentiate the postmeasurement samples from the as synthesized samples (30HAr), the prefix “5HAr” is used throughout the paper. Table 1 summarizes the composition and notation of the samples.

example, the authors in ref 23 claimed that europium enters the structure as Eu3+. Several studies demonstrated that EuTiO3 contains purely divalent Eu2+.1,22 Another important difference between EuTiO3 and SrTiO3 is the presence of a narrow Eu 4f band forming the valence band (VB) maximum in EuTiO3, while the top of the VB of SrTiO3 is formed by O 2p and Ti d bands.24 The Eu 4f electrons are not expected to contribute significantly to the chemical bonding,25 since they are localized and shielded by the Eu 5s and 5p electrons. Besides cationic- and anionic-substitutions, oxygen vacancies can also supply additional charge carriers. Oxygen vacancies act as electron dopants and strongly influence a number of material properties, such as electronic, magnetic,26 catalytic,27 and optical properties.28 Electronic transport properties of SrTiO3 strongly depend on the oxygen content.29−31 Oxygen-deficient SrTiO3 was the first oxide discovered to be superconducting.32 In contrast, the study of oxygen-deficient EuTiO 3 is accompanied with difficulties due its instability in air at elevated temperatures. At room temperature and under oxidative conditions, the pyrochlore phase Eu2Ti2O7 is the most thermodynamically stable configuration due to the propensity of Eu2+ to oxidize to Eu3+.33 However, below 723 K, the kinetics of this reaction is slow and no significant oxidation can be observed. Henderson et al.34 annealed EuTiO3 in an oxidative atmosphere and reported the transition to Eu2Ti2O7 through an amorphous phase. Hence, highly reductive conditions are required in order to create EuTiO335 and oxygen deficient EuTiO3‑δ. The aim of this work is twofold: (1) to investigate the effects of replacing Eu by the more abundant Sr in EuTiO3, and (2) to investigate the effects of oxygen vacancies on the electronic transport properties and structural properties. Five batches of SrxEu1−xTiO3‑δ powders with x = 0.00, 0.03, 0.25, 0.75, and 0.97 were produced in highly reducing atmosphere to create oxygendeficient samples. A second group of samples with slightly higher oxygen content was produced in a less reducing atmosphere. The differing oxygen content of the samples was determined by thermogravimetric analysis (TGA). We demonstrate that a slight variation in oxygen stoichiometry of SrxEu1−xTiO3‑δ samples affects the Seebeck coefficient, electrical resistivity, and optical absorption coefficient.

Table 1. Summary of the Composition and Notation of the Studied Samples Sr x x x x x

= = = = =

0.00 0.03 0.25 0.75 0.97

sample heated in 30% H2 in Ar

sample heated in 5% H2 in Ar

30HAr-ETO 30HAr-SETO3 30HAr-SETO25 30HAr-SETO75 30HAr-SETO97

5HAr-ETO 5HAr-SETO3 5HAr-SETO25 5HAr-SETO75 5HAr-SETO97

Powder XRD data of all samples were collected at room temperature using a PANanalytical X’Pert PRO system equipped with an X’Celerator linear detector operating in Bragg−Brentano geometry (θ/2θ) and a Johansson monochromator (Cu Kα1 radiation, 1.5406 Å). The samples were ground and packed into 16 mm diameter sample holders. The diffractometer was equipped with incident- and diffracted-beam Soller slits (0.4 rad). The diffraction patterns were recorded between 10° and 140° (2θ) at a continuous scan rate of 0.00483 deg/s. Peak profiles were determined by using the LeBail matching36 as implemented in the Fullprof program.37 Peak shape was described by a pseudo-Voigt function with additional parameters to account for peak asymmetry below 40° (2θ). Thermogravimetric analysis (TGA) was employed to assess the oxygen content of as-synthesized samples (30HAr) and after the high-temperature measurement (5HAr). The experiments were performed with a Netzsch STA 409 CD thermobalance. Around 150 mg of powder was oxidized by heating it up in air from 313 to 1523 K at a rate of 5 K/min. The oxygen nonstoichiometry parameter (δ) was evaluated from the weight gain during the oxidation experiment (eq 1). The errors were calculated from the instrumental error of the TGA instrument (0.05 mg). In a preliminary study with EuTiO3, the pyrochlore phase Eu2Ti2O7 did not show any observable weight change in a wide range of oxygen partial pressure pO2 (10−7 to 1 atm) at 1473 K. Therefore, we can treat the pyrochlore phase as oxygen stoichiometric and attribute the differences in weight gain during oxidation of EuTiO3‑δ to Eu2Ti2O7 solely to the different oxygen content within the perovskite phase.

2. EXPERIMENTAL METHODS A series of five oxide powders with general composition SrxEu1−xTiO3‑δ (x = 0.00, 0.03, 0.25, 0.75, 0.97) was prepared by solid-state reaction. Stoichiometric proportions of SrCO3 (≥98%; Sigma-Aldrich), Eu2O3 (99.9% purity; Metall Rare Earth Limited), and TiO2 (puriss.; Sigma-Aldrich) were mixed and ball milled for 3 h at 200 rpm for homogenization. The mixture was first calcined at 1123 K for 5 h in order to eliminate the carbonates. Then the powders were heated in a furnace (Carbolite, Type STF 15/450) specifically designed for high-temperature synthesis in hydrogen containing atmospheres. The samples were annealed at 1553 K for 10 h in a dynamic (570 mL/min) 30% H2/70% Ar atmosphere. The resulting samples after the synthesis will be denoted with the prefix “30HAr” in accordance with the synthesis gas conditions. The powders were pressed into disk-shaped pellets (10 mm diameter and 2 mm thick) using first 2 × 104 Pa uniaxial pressure, and then 1.8 × 108 Pa isostatic pressure. Afterward, the pellets were heated at 1553 K for 8 h under the same gas conditions (30HAr) for sintering and densification. The calculation of the densities (d) was based on the dimension

2EuTiO3‐δ + (δ + 0.5)O2 → Eu 2Ti 2O7

(1)

The sintered pellets were cut into small rectangular pieces of approximately 2 × 2 × 1 mm3 for X-ray photoelectron spectra (XPS) measurements. The samples were clamped onto a sample plate using molybdenum clips with 2 mm diameter apertures. X-ray photoelectron spectra were acquired on a Physical Electronics (PHI) Quantum 2000 photoelectron spectrometer using monochromatic Al Kα radiation and a hemispherical capacitor electron-energy analyzer equipped with 7822



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a channel plate and a position-sensitive detector. The area analyzed was typically of 150 μm diameter. Quantitative atomic concentrations were obtained from the different peak areas after Shirley-type background subtraction using the sensitivity factors given by the Multipack library. The electron takeoff angle was 45° and the analyzer operated in the constant pass energy mode. The pass energy used for the detail spectra of the C 1s, O 1s, Eu 3d, and Ti 2p core levels was 46.95 eV to yield a total analyzer energy resolution of 0.95 eV (for Ag 3d electrons). The samples were sputtered for 10 min with 1 kV Ar+ ions, and rastered over a 2 × 2 mm2 area to remove surface contamination. Approximately 20 nm was removed. To compensate for charging, the spectra of the unsputtered samples were shifted so that the C 1s signals coincided with the reference position for natural carbon contamination (as the samples were exposed to air), which is at 284.8 eV. The peak position of the O 1s signal of the unsputtered samples (529.5 eV) was taken as the calibration binding energy of the sputtered samples. UV−visible diffuse reflectance spectra (UV−vis DRS) were collected with a UV-3600 Shimadzu UV−vis-NIR spectrophotometer with an integrating sphere over a spectral range of 200−2500 nm (6.20−0.50 eV). The powders were pressed into disk-shaped pellets and glued to a quartz glass slide. The reflectance (R) data obtained were converted to the Kubelka− Munk absorbance F(R) using the following relation:38 F (R ) =

(1 − R )2 2R

Table 2. Summary of the X-ray Diffraction (XRD) LeBail Fitting Parametersa

a

sample

Rp

Rwp

χ2

a (Å)

30HAr-ETO 5HAr-ETO 30HAr-SETO3 5HAr-SETO3 30HAr-SETO25 5HAr-SETO25 (*) 30HAr-SETO50 (*) 5HAr-SETO50 30HAr-SETO75 5HAr-SETO75 30HAr-SETO97 5HAr-SETO97

5.56 5.53 5.32 5.22 5.5 5.44 6.68 6.26 7.66 7.73 9.17 8.54

7.11 7.05 6.77 6.64 6.98 6.92 8.43 7.96 9.83 9.8 12.8 11.9

1.37 1.33 1.27 1.24 1.24 1.21 1.55 1.35 1.55 1.55 1.92 1.65

3.9041(1) 3.9042(1) 3.9045(1) 3.9046(1) 3.9047(1) 3.9048(1) 3.9049(1) 3.9049(1) 3.9051(1) 3.9050(1) 3.9053(1) 3.9052(1)

The additional samples with x = 0.50 are indicated with asterisks.

(2)

The electrical resistivity (ρ) and Seebeck coefficient (S) measurements were performed on bar-shaped pellets of about 2 × 2 × 10 mm3 dimensions. The low-temperature (3 K < T < 300 K) measurements were performed with a Physical Property Measurement System (PPMS) (Quantum Design) in vacuum. ρ and S at high temperatures (300 K < T < 1289 K) were determined with a RZ200li system (Ozawa Science, Japan) in a reducing atmosphere of 5% H2 in Ar (5HAr) at 50 mL/min. The relative errors of the measured values of both ρ and S were evaluated to be 5%39 and the error bars were added accordingly.

Figure 1. Lattice parameter (a) as a function of the Sr content for 30HAr and 5HAr samples. The additional samples with x = 0.50 are indicated with an asterisk.

Eu-rich samples have slightly larger lattice parameters compared to their 30HAr counterparts (Figure 1). In order to confirm this trend, one more sample, with x = 0.50, was synthesized and measured. The results of the XRD fittings are shown in Table 2 and Figure 1 (the additional samples are indicated with asterisks). The lattice expansion of the Sr-rich samples, exposed to strongly reducing conditions (30HAr), is attributed to the presence of Ti3+ on the B-site as r(Ti3+) = 0.810 Å and r(Ti4+) = 0.745 Å, and is confirmed by XPS measurements (explained below). However, the lattice parameter expansion observed in Eu-rich samples exposed to less reducing conditions (5HAr) is possibly due to the partial substitution of Eu onto the B-site. Partial oxidation of Eu2+ to Eu3+ might facilitate Eu3+ occupation of the B-site. This has been observed in other titanates such as SrTiO341 and CaTiO3,42,43 and it is more likely to occur for small A-site cations.44 This effect is probably stronger in the Eu-rich samples, since Sr2+ cannot be oxidized to a higher oxidation state. The lattice parameter of the sample with x = 0.50, midway in the Sr−Eu series, exposed to strongly reducing conditions is almost unchanged after exposure to less reducing conditions. In this composition, both expansion and contraction effects are roughly equally counteracting, thereby supporting the proposed trend.

3. RESULTS AND DISCUSSION The XRD patterns of both the 30HAr and 5HAr series SrxEu1−xTiO3‑δ samples were found to be phase pure and can be indexed to the cubic Pm3̅m space group. The maximum tolerable oxygen nonstoichiometry was not reached in the samples as no traces of the pyrochlore Eu2Ti2O7 phase or other secondary phases were found. The lattice parameters (a) of the measured samples were calculated by LeBail profile fitting. They increased systematically with Sr content, from a = 3.9041(1) Å for x = 0.00 to a = 3.9053(1) Å for x = 0.97 (30HAr) (Table 2 and Figure 1). These values are in reasonable agreement with the lattice parameters reported in literature for EuTiO3 and SrTiO3, 3.90416 and 3.905 Å,17 respectively. The reason for the unit cell expansion with increasing Sr content has been reported in literature.11 It is attributed to the lower electronegativity and the slightly bigger ionic radius of Sr2+ compared to Eu2+ [χ(Sr) = 0.95, χ(Eu) = 1.2 and r(Sr) = 1.36 Å, r(Eu) = 1.35 Å].40 In addition to the anticipated lattice parameter expansion, a comparison between the 30HAr and 5HAr series lattice parameters reveals another trend: The 30HAr Sr-rich lattice parameters are slightly larger compared to their 5HAr counterparts; conversely, the 5HAr 7823



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(Sr,Eu)TiO3‑δ perovskite phases. Especially in the case of the 30HAr samples, the difference between the calculated δ values and the error of the measurements is very small and, therefore, near to the stoichiometric oxygen value. Nevertheless, it is clear from the results that the 30HAr samples have a slightly lower oxygen stoichiometry than the 5HAr-ETO samples, confirming the influence of the two atmospheres. It should be mentioned that the oxygen uptake in the 5HAr samples occurs because of residual oxygen partial pressure present in the annealing chamber. This is likely a consequence of water condensed on the surfaces of the sample and/or on the chamber, or of small leakages. Moreover, it illustrates the strong propensity for Eu to oxidize to Eu3+. X-ray photoemission spectra of the Ti 2p and Eu 3d core levels were studied to analyze the oxidation states of the cations. Determination of the Ti oxidation state by XPS is a critical issue since during Ar-ion sputtering a preferential loss of oxygen is generated resulting in Ti3+ and Ti2+ states.45 Therefore, the valence state of Ti was determined from the XPS analysis of the unsputtered samples. It has to be considered that air-exposed Ti compounds have a 1−2 nm thick surface layer of TiO2. Since the Ti 2p electrons acquired are partially from the natural surface oxide layer and partially from below, at least a qualitative determination of the bulk Ti valence states can be measured by XPS surface analysis. Figure 3a shows the peak fitting of the Ti 2p doublet of the 30HArSETO3 and 5HAr-SETO3 samples, as representative for the group of 30HAr and 5HAr samples, respectively. The main peaks are centered at 458.6 eV (2p3/2) and 464.3 eV (2p1/2), and correspond to Ti4+.46 The shoulders observed in the spectra of the 30HAr sample originate from Ti3+ (upper panel of Figure 3a).46 The deconvoluted Ti3+ peaks are located at 457.4 eV (2p3/2) and 463.0 eV (2p1/2). The presence of trivalent Ti3+ ions in the 30HAr samples is a result of the charge compensation necessary to balance the oxygen vacancies. The peaks of the 5HAr sample are symmetric since they originate uniquely from Ti4+ and contributions of other oxidation states of Ti are not present [lower panel of Figure 3a]. This result supports the existence of oxygen vacancies in the strongly reduced samples (30HAr) and not in the 5HAr samples; and it correlates to the larger unit cell of the 30HAr Sr-rich samples found by XRD [r(Ti3+) > r(Ti4+)]. The photoelectron spectra of Eu 3d were measured after cleaning part of the surface oxide layer by sputtering. It is important to emphasize the strong sensitivity of surface Eu2+ to oxidize to Eu3+ in contact with air. Orlowski et al.47 collected the Eu 3d photoemission spectra of a Eu2+-based compound subjected to Ar ion sputtering and they found Eu3+ traces on the surface of the sample after sputtering. Accordingly, we also expect to record Eu3+ signals of the remaining oxidized surface layer. Figure 3b shows the spectra of 30HAr-SETO3 and 5HArSETO3 samples. The main peaks appear at 1124.4 eV (3d5/2) and 1154.0 eV (3d3/2), and correspond to the Eu2+ doublet, while the lower intensity doublets located at higher binding energies correspond to trivalent Eu3+ [at 1133.1 eV (3d5/2) and 1161.0 eV (3d3/2)].48 There is a contribution from satellite excitations, which overlap the Eu3+ main peak47 and is indicated with an asterisk in Figure 3b. It is important to note that a large part of the Eu3+ peaks originates from surface oxidation as reported in ref 47. Consequently, it is only possible to do a qualitative and comparative analysis of the spectra. We calculated the Eu3+/Eu2+ ratios of 5HAr-SETO3 and 30HArSETO3, resulting in 0.27 and 0.21, respectively, confirming the

The TGA oxidation measurements demonstrated that the different reducing power of the annealing atmospheres (30HAr and 5HAr) resulted in samples with slightly different oxygen stoichiometries. Here, we show the oxidation curves of the unsubstituted samples 30HAr-ETO and 5HAr-ETO as representative for the SrxEu1−xTiO3‑δ solid solution (Figure 2). Additionally, the mass gain values (Δm) and oxygen

Figure 2. Mass gain of 30HAr-ETO and 5HAr-ETO during the oxidation in air.

nonstoichiometry parameters (δ) of all the samples are listed in Table 3. Figure 2 shows that oxygen uptake in both samples Table 3. Mass Gain (Δm) and Oxygen Nonstoichiometry Parameter (δ) of SrxEu1−xTiO3‑δ from Oxidative Thermogravimetric Analysis (TGA) sample

mass gain Δm (%)

oxygen nonstoichiometry parameter δ

stoichiometry

30HAr-ETO 5HAr-ETO 30HAr-SETO3 5HAr-SETO3 30HAr-SETO25 5HAr-SETO25 30HAr-SETO75 5HAr-SETO75 30HAr-SETO97 5HAr-SETO97

3.27(3) 3.14(2) 3.06(3) 3.03(3) 2.57(3) 2.47(3) 0.95(3) 0.85(4) 0.03(3) −0.08(3)

0.007(5) −0.013(3) −0.014(5) −0.019(5) −0.003(4) −0.017(5) −0.006(4) −0.019(5) −0.012(4) −0.024(4)

EuTiO2.993(5) EuTiO3.013(3) Sr0.03Eu0.97TiO3.014(5) Sr0.03Eu0.97TiO3.019(5) Sr0.25Eu0.75TiO3.003(4) Sr0.25Eu0.75TiO3.017(5) Sr0.75Eu0.25TiO3.006(4) Sr0.75Eu0.25TiO3.019(5) Sr0.97Eu0.03TiO3.012(4) Sr0.97Eu0.03TiO3.024(4)

starts at around 600 K. The mass gain of 30HAr-ETO [Δm = 3.27(3)%] was slightly higher than that of 5HAr-ETO [Δm = 3.14(2)%] indicating a lower oxygen content in the highly reduced sample. The molar masses of the initial powders [M(EuTiO3‑δ)] were calculated with eq 3 and δ parameters were derived from eq 4. 2M(EuTiO3‐δ ) =

M(Eu 2Ti 2O7 ) 100% + Δm (%)

(3)

2M(EuTiO3‐δ ) = 2[M(Eu) + M(Ti) + (3 − δ)M(O)] (4)

Table 3 shows that the δ values of all the SrxEu1−xTiO3‑δ are very small compared with other perovskite oxides,29 indicating the low tolerance for deviations in oxygen stoichiometry in the 7824



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Figure 3. (a) Deconvoluted X-ray photoelectron Ti 2p spectra of the 30HAr-SETO3 and 5HAr-SEO3 samples. Ti4+ is the main oxidation state, although 30HAr-SETO3 has a small contribution of Ti3+. (b) Eu 3d spectra of 30HAr-SETO3 and 5HAr-SETO3. The asterisks show the satellite excitations. Only one sample of each series is shown, the results presented here are representative for all samples measured.

Figure 4. (a−e) Tauc plots of the samples and band gap (Eg) values. (f) Summary plot of Eg versus the Sr substitution (x).

coefficients [F(R)] were plotted using Tauc’s relation (eq 5)49 in order to obtain the Eg values.

slightly higher amount of Eu3+ in the former. The other samples with x = 0.00, 0.25, 0.75, and 0.97 (not shown here) also showed higher Eu3+/Eu2+ for the 5HAr sample than for the corresponding 30HAr sample. This result supports the higher Eu3+ content suggested for the explanation of the larger a parameter of the Eu-rich 5HAr samples. The optical band gaps (Eg) were obtained from UV−vis spectroscopy in order to explore the effect of Sr and oxygen content variations. The Kubelka−Munk optical absorption

F(R )hν ∝ (hν − Eg )b

(5)

The product hν represents the energy of the incident photon and the exponent b depends on the type of transition: b = 2 for indirect band gaps and b = 1/2 for direct band gaps. Thus, the plot (F(R)hν)1/b vs hν follows a linear dependence in the region of the valence band to conduction band (VB-CB) transition, and the linear extrapolation to the abscissa gives the 7825



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Figure 5. Spectral dependence of lnF(R) in the region near the band edge. The three different photoenergy regions described by Urbach are shown in (e).

Eg value. Figure 4 shows the Tauc plots of all the samples. The absorption features of samples with 0.03 ≤ x ≤ 0.75 are very similar to the unsubstituted EuTiO3 and they have direct band gaps near 1 eV (from 0.96 to 1.14 eV). On the contrary, the band gaps of x = 0.97 are considered indirect due to their similarity with SrTiO3 (Eg = 3.25 eV)50 and their values are ∼3.1 eV. However, a small feature near 1 eV can still be seen in the spectra of x = 0.97 coming from few transitions between Eu 4f electrons and the CB (not shown here). Figure 4f summarizes all the Eg extracted from Tauc plots. The values of 30HAr samples are slightly lower than those of the corresponding 5HAr samples probably due to the band gap closure or tailing of the band edge induced by the oxygen vacancies. The calculated band gaps of EuTiO3 (Eg = 0.96 eV and Eg = 0.98 eV for 30HAr-ETO and 5HAr-ETO samples, respectively) are in agreement with the value 0.93 ± 0.07 eV reported by Lee et al.35 The sample with x = 0.03 shows the same values as EuTiO3. Across the samples with 0.25 ≤ x ≤ 0.97 the band gap increases with increasing Sr substitution. An abrupt increase occurs in the x = 0.97 sample. It was demonstrated by Ranjan et al.51 that the electronic density of states (DOS) of the top of the VB of EuTiO3 is mostly formed by the Eu 4f orbitals, separated by ∼1 eV from the CB. However, in SrTiO3, the most prominent contribution to the top of the VB originates from O 2p levels, which is separated from the CB band by ∼3.25 eV.50 Our results show the relevance of the Eu 4f states at the top of the VB of SrxEu1−xTiO3‑δ since all the spectra of the samples with x ≤ 0.75 are dominated by the charge transfer between the Eu 4f band and the CB.

In disordered semiconductors, localized states extended in the band gap can appear and are called tail states.52 The optical absorption coefficient near the band edge can usually be separated in three different photoenergy regions as plotted in Figure 5e. (1) At energies above the band gap, the absorption coefficient follows Tauc’s relation, as described in eq 5. (2) Directly below, the absorption follows the empirical Urbach rule53 and varies exponentially with the incident photon energy as described by eq 6: ⎛ hν ⎞ F(R ) = F0 exp⎜ ⎟ ⎝ EU ⎠

(6)

where F0 is a constant and EU is the Urbach energy. EU denotes the energy width of the localized states in the band gap (tails states). (3) At even lower energies, the absorption is controlled by transitions from tail to tail states.54 In Figure 5, the ln F(R) vs hν curves are plotted. The linear dependence described by Urbach is only clearly visible in the samples with highest Sr content (x = 0.97) (Figure 5e), and the values of the Urbach energies can be extracted from the slope of the linear part of the plot (EU = 156 meV). The EU value can be interpreted as the width of the Urbach tail. In this case, the calculated EU values for the 30HAR and 5HAr samples are the same. The samples with 0.00 ≤ x ≤ 0.75 do not clearly show the three photoenergy regions and the linear Urbach dependence, probably due to their small band gap relative to the defect states. All the 30HAr samples present higher absorption below the Urbach tail than the corresponding 5HAr samples, indicating a higher amount of tail to tail transitions. This result confirms the higher presence 7826



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Figure 6. Electrical resistivity (ρ) and Seebeck coefficient (S) of SrxEu1−xTiO3‑δ samples with (a) x = 0.00, (b) x = 0.03, (c) x = 0.25, (d) x = 0.75, and (e) x = 0.97. Region I: low T measurement of the 30HAr samples. Region II: in situ oxygen uptake where the 5HAr samples are formed. Regions III and IV: high and low T measurement of the 5HAr samples.

and the cooling slopes strongly differ from each other due to in situ oxygen uptake during the measurements. Each graph is divided into four regions: region I refers to the low-temperature measurements of the 30HAr (oxygen-deficient) samples; region II denotes the high-temperature measurements in 5HAr atmosphere where in situ oxygen uptake occurs, forming the

of optically active defects due to the oxygen vacancies in the 30HAr samples. The electrical resistivity and Seebeck coefficient of the SrxEu1−xTiO3‑δ samples were measured as a function of temperature in the range of 3 K < T < 1289 K in a heating and cooling cycle (Figure 6). The data points of the heating 7827



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Figure 7. Linear variation of ln ρ as a function of T−1/4 in the low-temperature regime of SrxEu1−xTiO3‑δ samples with (a) x = 0.00, (b) x = 0.03, (c) x = 0.25, and (d) x = 0.75.

the Ti ions deviates Ti4+ to intermediate values Ti3+/4+ (as proved by XPS) to maintain the overall charge, and the electrical resistivity decreases significantly. This effect has been named valence control by Fujimori.55 Samples with 0 ≤ x ≤ 0.75 showed semiconducting behavior (dρ/dT < 0) in the whole temperature range for both oxygen stoichiometries (30HAr and 5HAr) (Figure 6a−d). However, the sample with x = 0.97 has metallic behavior in the temperature range of 156 K < T < 895 K (regions I and II) and there is a transition from metallic to semiconducting behavior at T = 895 K (region II) (Figure 6e). The low-temperature ρ of samples 0 ≤ x ≤ 0.75 (region I and IV) can be partially described by the three-dimensional Mott variable range hopping (VRH):56

5HAr samples; region III corresponds to the high-temperature measurements recorded upon cooling of the 5HAr samples; and finally region IV refers to the low-temperature measurements of the 5HAr samples. A dotted line situated at 300 K divides the graphs into the low- and the high-temperature regions, and the arrows indicate the order in which the measurements were performed (Figure 6). The temperature of the 30HAr to 5HAr oxidation can be approximated from the region where the heating and cooling slopes begin to coincide (indicated with asterisks). These temperatures are around 590 K for 0 ≤ x ≤ 0.25 (confirmed by TGA); 780 K for x = 0.75, and 890 K for x = 0.97. The increase in oxidation temperature with Sr content is due to the greater electron affinity of Ti3+ compared to Eu2+ as the oxidation of Ti3+ to Ti4+ and Eu2+ to Eu3+ is believed to dominate the oxidation behavior of the Srrich and the Eu-rich samples, respectively. The low-temperature ρ of the 30HAr samples (region I) is several orders of magnitude lower than that of the 5HAr samples (region IV): for samples with x = 0.00 and 0.97, ρ changes 5 orders of magnitude; for x = 0.03 the change reaches up to 4 orders of magnitude, and for x = 0.25 and 0.75 there is a difference of 3 orders of magnitude. Such a decrease of ρ clearly demonstrates the strong effect of electron doping caused by the oxygen vacancies in the 30HAr samples: The mean valence of

⎛ T ⎞1/4 ρ = ρ0 exp⎜ 0 ⎟ ⎝T ⎠

(7)

3

where T0 = 16α /kBN(EF) is a constant which can be calculated from the slope of the plot of ln ρ vs T−1/4, α is the electron wave function decay constant and N(EF) is the density of states at the Fermi level. Figure 7 shows the linear behavior of ln ρ vs T−1/4 in the indicated temperature ranges. Although Mott described the T−1/4 behavior for very low temperatures, Pal et al.57 and other authors58,59 successfully applied the VRH conduction at temperatures above 200 K. The slope of ln ρ vs 7828



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Table 4. Slopes of the Electrical Resistivity Fits at Low and High Temperatures region

slope

x = 0.00 (ETO)

x = 0.03 (SETO3)

x = 0.25 (SETO25)

x = 0.75 (SETO75)

reg. I (30HAr) reg. IV (5HAr) reg. III (5HAr)

(T0)1/4 [(K)1/4] (T0)1/4 [(K)1/4] EA (meV)

19.60 146.39 321.15

14.56 142.48 347.54

39.36 155.74 349.54

16.89

T−1/4 of the 30HAr samples is lower compared to the corresponding 5HAr samples (for 0 ≤ x ≤ 0.25), indicating the presence of higher density of states at the Fermi level N(EF) in the reduced samples57 (Table 4). This result can be correlated with the higher amount of defect states observed in the UV−vis measurements of the 30HAr samples. The electrical resistivity of 5HAr-SETO75 increases quickly with decreasing temperature, therefore ρ and S were out of the measurement range below 300 K. The low-temperature resistivity of SETO97 could not be fitted to Mott−VRH behavior. Above room temperature, electrical conduction of all the samples occurs via thermal excitation of the electrons into the conduction band (CB), and ρ is proportional to the exponential term exp(EA/kBT), where EA is the activation energy. The linear fit of lnρ against 1/T provides the EA values (Table 4). Samples with 0.03 ≤ x ≤ 0.75 behave similarly to EuTiO3 (Figure 6a− d). The lowest values of ρ are reached at the highest measured temperatures due to their semiconducting behavior (ρ ≈ 4 × 10−4 Ω·m at 1287 K for 0.00 ≤ x ≤ 0.25 and ρ ≈ 9 × 10−4 Ω·m at 1288 K for x = 0.75). The lowest resistivity value of the Srrich sample (x = 0.97), ρ ≈ 5 × 10−4 Ω·m, was reached at low temperature (156 K) in the metallic regime (region I). No significant dependence of the electrical resistivity upon the Sr content was detected for the samples with 0.00 ≤ x ≤ 0.75, similar to the case of (Sr,Eu)Ti0.8Nb0.2O3 reported by Kato et al.14 These results demonstrate that the electrical conductivity of SrxEu1−xTiO3‑δ cubic perovskites with x up to x = 0.75 is dominated by the Eu 4f electrons and reflects the charge transfer between the occupied Eu 4f and the conduction band, as shown by UV−vis spectroscopy. On the other hand, the ρ behavior of the sample with highest Sr content, x = 0.97, deviates significantly from the ρ behavior of EuTiO3: 30HArSETO97 is metallic at 156 K < T < 895 K and 5HAr-SETO97 becomes less resistive at low temperatures compared with the other 5HAr samples. The Seebeck coefficients of SrxEu1−xTiO3‑δ samples during the heating and cooling cycle are shown in Figure 6. The Seebeck coefficients were negative throughout the entire temperature range indicating that electrons are the dominant mobile charge carriers. The S values of SrxEu1−xTiO3‑δ are particularly high compared to other semiconducting perovskitetype oxides.60−62 In most of the T range, the 30HAr annealed samples (region I) show lower S values than the 5HAr samples due to the higher charge carrier density arising from the oxygen vacancies. The 5HAr samples reach large negative values (∼ − 1000 μV/K) at about 250−310 K for samples with 0.00 ≤ x ≤ 0.75 and at 580 K for x = 0.97, similarly as it was shown in ref 12 for EuTiO3. The change in the sign of the S(T) slope at the mentioned temperatures does not correspond to the typical semiconducting behavior (d|S|/dT < 0) which is observed in all the samples (except for x = 0.97 that is metallic at 156 K < T < 895 K). The origin of the pronounced S(T) peaks might be attributed to the hybridization of the localized 4f orbitals of Eu ([Xe] 4f7 6s2) with the partially filled Ti 3d states. This hybridization has been described by several authors22,25 and is

416.24

x = 0.97 (SETO97)

255.88

considered to be the dominant cause of the strong spin−lattice coupling in EuTiO3. Bauer et al.63 described that the strong hybridization between localized 4f states and delocalized conduction electrons results in the mixing of the respective electrons. This mixing destabilizes the local 4f moments and causes scattering resonances near the Fermi level, leading to an energy pseudogap at the Fermi level, which enhances the thermopower (≡ S). In the case of EuTiO3, a similar phenomenon of destabilization of the local 4f moments through hybridization might occur as a consequence of the lattice instabilities due to Eu atomic displacements below room temperature, described by Bessas et al.22 This mechanism can possibly explain the high S(T) peak of SrxEu1−xTiO3‑δ. Indeed, the sample with lowest Eu amount (x = 0.97) shows a smoother peak compared to the samples with higher Eu content. In ref 64, a similar kind of thermopower behavior was reported for Ce intermetallics under high pressures. The hybridization strength (Γ) between the Ce 4f electrons and the conduction band increases with pressure. When Γ is higher than 2 times the crystal field splitting (Γ > 2Δ), the Seebeck curve presents a single peak at higher temperatures (∼300−600 K). Such behavior is in qualitative agreement with the Seebeck coefficient of valence f luctuators.64,65 In our case, the peak of 5HAr samples is more pronounced compared to the 30HAr samples. This might be due to valence fluctuations of Eu (mixed valence state Eu2+/3+ as verified by XPS). Another possibility to explain the higher intensity of 5HAr Seebeck coefficient peaks is the Eu−O−Ti−O−Eu superexchange, which according to Akamatsu et al.3 cannot be completely ruled out. The Seebeck coefficients of the 0.03 ≤ x ≤ 0.75 samples are similar to EuTiO3, whereas the sample with x = 0.97 undergoes a change in the S(T) slope in region (II) due to the metal− semiconductor transition. In region III, there is a slope change analogous to the samples with higher Eu content. Again, the Eu 4f electrons dominate the electron transport properties in all samples with x ≤ 0.75. The Eu 4f electrons in EuTiO3 are localized and shielded by the Eu 5s and 5p electrons and are, therefore, not expected to contribute significantly to the chemical bonding;25 nevertheless, the Eu 4f band strongly influences the electronic transport properties of SrxEu1−xTiO3‑δ samples.

4. CONCLUSIONS Perovskite-type SrxEu1−xTiO3‑δ (x = 0.00, 0.03, 0.25, 0.75, 0.97) samples were treated in highly reducing atmosphere (30HAr) to create oxygen vacancies. The oxygen content was slightly increased during subsequent high-temperature measurements of ρ and S in less reducing atmosphere (5HAr). The 30HAr samples remained cubic after being oxidized to the 5HAr samples. Further, the presence of higher amounts of active defect states in the 30HAr samples was confirmed by UV−vis measurements and ln F(R) vs hν plots. ρ and S measurements reflected the strong influence of the varied oxygen content on the transport properties. The 30HAr samples exhibited lower electrical resistivity (up to 5 orders of magnitude lower than the 7829



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(11) Muta, H.; Ieda, A.; Kurosaki, K.; Yamanaka, S. Thermoelectric Properties of Lanthanum-Doped Europium Titanate. Mater. Trans. 2005, 46, 1466−1469. (12) Sagarna, L.; Shkabko, A.; Populoh, S.; Karvonen, L.; Weidenkaff, A. Electronic Structure and Thermoelectric Properties of Nanostructured EuTi1−xNbxO3−δ (x = 0.00; 0.02). Appl. Phys. Lett. 2012, 101, 033908. (13) Sagarna, L.; Rushchanskii, K. Z.; Maegli, A.; Yoon, S.; Populoh, S.; Shkabko, A.; Pokrant, S.; Ležaić, M.; Waser, R.; Weidenkaff, A. Structure and Thermoelectric Properties of EuTi(O,N)3±δ. J. Appl. Phys. 2013, 114, 033701. (14) Kato, K.; Yamamoto, M.; Ohta, S.; Muta, H.; Kurosaki, K.; Yamanaka, S.; Iwasaki, H.; Ohta, H.; Koumoto, K. The Effect of Eu Substitution on Thermoelectric Properties of SrTi0.8Nb0.2O3. J. Appl. Phys. 2007, 102, 116107. (15) Guguchia, Z.; Shengelaya, A.; Keller, H.; Köhler, J.; BussmannHolder, A. Tuning the Structural Instability of SrTiO3 by Eu Doping: The Phase Diagram of Sr1−xEuxTiO3. Phys. Rev. B 2012, 85, 134113. (16) Janes, D. L.; Bodnar, R. E.; Taylor, A. L. Europium Barium titanate - A Magnetic Ferroelectric Compound. J. Appl. Phys. 2008, 49, 1452−1454. (17) Chen, Z.-X.; Chen, Y.; Jiang, Y.-S. Comparative Study of ABO3 Perovskite Compounds. 1. ATiO3 (A = Ca, Sr, Ba, and Pb) Perovskites. J. Phys. Chem. B 2002, 106, 9986−9992. (18) Fleury, P. A.; Scott, J. F.; Worlock, J. M. Soft Phonon Modes and the 110 K Phase Transition in SrTiO3. Phys. Rev. Lett. 1968, 21, 16. (19) Bussmann-Holder, A.; Köhler, J.; Kremer, R.; Law, J. Relation between Structural Instabilities in EuTiO3 and SrTiO3. Phys. Rev. B 2011, 83, 212102. (20) Allieta, M.; Scavini, M.; Spalek, L. J.; Scagnoli, V.; Walker, H. C.; Panagopoulos, C.; Saxena, S. S.; Katsufuji, T.; Mazzoli, C. Role of Intrinsic Disorder in the Structural Phase Transition of Magnetoelectric EuTiO3. Phys. Rev. B 2012, 85, 184107. (21) Kim, J.-W.; Thompson, P.; Brown, S.; Normile, P. S.; Schlueter, J. A.; Shkabko, A.; Weidenkaff, A.; Ryan, P. J. Emergent Superstructural Dynamic Order due to Competing Antiferroelectric and Antiferrodistortive Instabilities in Bulk EuTiO3. Phys. Rev. Lett. 2013, 110, 027201. (22) Bessas, D.; Rushchanskii, K. Z.; Kachlik, M.; Disch, S.; Gourdon, O.; Bednarcik, J.; Maca, K.; Sergueev, I.; Kamba, S.; Ležaić, M.; et al. Lattice Instabilities in Bulk EuTiO3. Phys. Rev. B 2013, 88, 144308. (23) Fajardo-Peralta, A.; Romero-Vargas, F. N.; Rosas-Bonilla, J. C.; Torres-Mayorga, J. L.; Mata, J.; Siqueiros, J. M. Characterization of Polycrystals of STO Doped with Europium: Sr1−xEuxTiO3. Integr. Ferroelectr. 2008, 101, 114−120. (24) Wei, W.; Dai, Y.; Guo, M.; Yu, L.; Huang, B. Density Functional Characterization of the Electronic Structure and Optical Properties of N-Doped, La-Doped, and N/La-Codoped SrTiO3. J. Phys. Chem. C 2009, 113, 15046−15050. (25) Birol, T.; Fennie, C. J. Origin of Giant Spin-Lattice Coupling and the Suppression of Ferroelectricity in EuTiO3 from First Principles. Phys. Rev. B 2013, 88, 094103. (26) Fang, D. Q.; Rosa, A. L.; Zhang, R. Q.; Frauenheim, T. Theoretical Exploration of the Structural, Electronic, and Magnetic Properties of ZnO Nanotubes with Vacancies, Antisites, and Nitrogen Substitutional Defects. J. Phys. Chem. C 2010, 114, 5760−5766. (27) Dai, X. P.; Wu, Q.; Li, R. J.; Yu, C. C.; Hao, Z. P. Hydrogen Production from a Combination of the Water−Gas Shift and Redox Cycle Process of Methane Partial Oxidation via Lattice Oxygen over LaFeO3 Perovskite Catalyst. J. Phys. Chem. B 2006, 110, 25856− 25862. (28) Cong, Y.; Li, B.; Yue, S.; Fan, D.; Wang, X. Effect of Oxygen Vacancy on Phase Transition and Photoluminescence Properties of Nanocrystalline Zirconia Synthesized by the One-Pot Reaction. J. Phys. Chem. C 2009, 113, 13974−13978. (29) Gong, W.; Yun, H.; Ning, Y. B.; Greedan, J. E.; Datars, W. R.; Stager, C. V. Oxygen-Deficient SrTiO3−x, x = 0.28, 0.17, and 0.08.

5HAr samples) due to the effect of the oxygen vacancies and the associated mixed Ti4+/3+ valence states (as shown by XPS). All of the samples with 0.00 ≤ x ≤ 0.75 are semiconducting with behavior similar to that of EuTiO3, due to the localized Eu 4f orbitals acting as dominant band of the electronic VB structure. The sample with the highest Sr content (30HAr-97) acts like a metal in the 156 K < T < 894 K temperature range. The Seebeck coefficients of each of the 5HAr samples peak near room temperature; the sample containing the least Eu (x = 0.97) possesses the least pronounced peak, however. The anomalous behavior of S(T) originates from the strong hybridization between the localized Eu 4f states and the delocalized Ti 3d electrons.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: weidenkaff@imw.uni-stuttgart.de. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

This research was supported by SNF (Swiss National Science Foundation) within the National Centre of Competence in Research (NCCR) “MaNEP − Materials with Novel Electronic Properties” and within the Sinergia programme (Grant No: CRSII2_136299/1 − Thermoelectric oxides TEO).

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Influence of the Oxygen Content on the Electronic Transport Properties

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