Oxygen Nonstoichiometry and Defect Chemistry of Perovskite

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Oxygen Nonstoichiometry and Defect Chemistry of PerovskiteStructured BaxSr1−xTi1−yFeyO3−y/2+δ Solid Solutions Melanie Kuhn,*,† Jae Jin Kim,† Sean R. Bishop,†,‡ and Harry L. Tuller†,‡ †

Department of Materials Science and Engineering, Massachusetts Institute of Technology, 77 Massachusetts Avenue, Cambridge, Massachusetts 02139, United States ‡ International Institute for Carbon Neutral Energy Research (WPI-I2CNER), Kyushu University, Nishi-ku Fukuoka, 819-0395, Japan ABSTRACT: The oxygen nonstoichiometry of mixed conducting perovskite-structured BaxSr1−xTi1−yFeyO3−y/2+δ (BSTF) (x = 0, 0.1, 0.5 and y = 0.05, 0.35) was measured by means of thermogravimetry as a function of oxygen partial pressure, pO2, in a temperature range typical for solid oxide fuel cell (SOFC) cathode applications. With increasing Ba content, the nonstoichiometry curve was shifted to higher pO2, indicating enhanced reducibility. Based on a defect chemical analysis of thermogravimetry derived nonstoichiometry data, the substitution of Ba for Sr on the A-site of STF was found to result in a decrease in both the reduction enthalpy and band gap energy, consistent with expectations that substitution of Sr by the larger Ba leads to a reduction in bond strength. A consequent increase in minority electron density and oxygen vacancy concentration are expected to result in enhancements in oxygen surface exchange kinetics and diffusivity and thereby cathode performance. The nonstoichiometry data obtained in this study also brought to light a previous underestimation of the minority electron density, by approximately a factor of 103, a key parameter believed to impact cathodic performance. KEYWORDS: BSTF, cathode, SOFC, defect equilibria

1. INTRODUCTION One of the primary aims of current research on solid oxide fuel cells (SOFC) focuses on the reduction of operating temperatures, implying restrictions on the choice of materials, especially on the cathode side.1−3 Kinetics of the already sluggish oxygen reduction reaction are negatively affected by reduced temperatures, and as a consequence, mixed-conducting oxide materials with high surface oxygen exchange and diffusion rates are necessary.4 SrTi1−xFexO3−x/2+δ (STF), a mixed conducting perovskite-structured oxide, is a candidate material both for application as an oxygen sensor and as an SOFC cathode.5−11 As a result, the conductivity of STF has been investigated, in detail, as functions of temperature, oxygen partial pressure (pO2), and iron content.5,7,10 The cathode performance of STF, in the form of thin films, was subsequently evaluated, with the surface oxygen exchange rate, kex, found to be the rate limiting step.6,12,13 This rate was found to depend on the minority charge carrier density, n, which, in turn, controls the charge transfer rate of electrons from electrode surface to adsorbed oxygen molecules.6 Following nondegenerate semiconductor statistics, the electron charge carrier density can be correlated to the position of the Fermi energy EF relative to the bottom of the conduction band Ec, by the following relationship: ⎛ E − EF ⎞ n = NC exp⎜ − C ⎟ kBT ⎠ ⎝ © 2013 American Chemical Society

with Nc the effective density of states in the conduction band, kB the Boltzmann constant, and T the temperature in K. An enhancement of kex could then be expected from either (i) a decrease in enthalpy for reduction, ΔHred, thereby lifting EF upward toward Ec, or by (ii) a reduction in the band gap, Eg, thereby lowering Ec. In the present work, we study the effect of substituting the larger cation, Ba (radius of 1.61 Å),14 onto the Sr (radius of 1.44 Å)14 A-site, on the key parameters ΔHred and Eg. With an increase in lattice constant, upon substitution of Ba for Sr, a reduction in bond strength is expected and thereby reduced ΔHred and Eg. To test this hypothesis, the oxygen nonstoichiometry of BaxSr1−xTi1−yFeyO3−y/2+δ (BSTF) is measured by means of thermogravimetric analysis (TGA) and values for ΔHred and Eg are obtained as a function of Ba doping by fitting the data to a defect chemical model. Previously, the authors investigated two compositions of BSTF and found behavior consistent with the above discussion.15 In the present work, four compositions of BSTF as well as one composition of STF are more extensively studied. While a defect model for STF was previously published by the authors, it was based largely on electrical measurements.7 As demonstrated below, the new more extensive data provided in this study requires a Received: February 16, 2013 Revised: June 24, 2013 Published: June 27, 2013

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dx.doi.org/10.1021/cm400546z | Chem. Mater. 2013, 25, 2970−2975

Chemistry of Materials

Article

refinement of the model, particularly with respect to the density of electrons at given conditions of temperature and pO2.

The condition for charge neutrality is given by // 2[V •• O ] + p = 2[Oi ] + n

2. THEORY A detailed defect chemical model, based on previous studies of the transport properties of STF,7 is summarized here. The +3 stoichiometric composition is defined as SrTi+4 1−xFex O3−x/2, where neither Fe nor Ti are considered as defect species and where the preferred oxidation state of Fe is +3, resulting in an inherent oxygen deficient perovskite.7 The intrinsic disorder is thus anion Frenkel disorder, where oxygen vacancies and interstitials (VO•• and Oi//, respectively, in Krö ger−Vink notation)16 are thermally generated by the following reaction: // OOx + Vix ↔ V •• O + Oi

assuming that cation defects are negligible below about 1000 °C.7 Lastly, the TGA measurement records the nonstoichiometry, given by δ=

// [V •• O ][Oi ] x [OO][V ix]

(13)

3. EXPERIMENTAL DETAILS

(2)

Single-phase BaxSr1−xTi1−yFeyO3−y/2+δ (BSTF) powders (x = 0, 0.1, 0.5 and y = 0.05, 0.35) were synthesized by the Pechini method.15 First, Ti[OCH(CH3)2]4 was dripped into ethylene glycol, followed by adding citric acid to form a metal−organic complex. After adding Ba(NO3)2, Sr(NO3)2, Fe(NO3)3·9H2O, and distilled water in the respective stoichiometric ratio, the resultant mixtures were stirred at 90 °C until a dry gel formed through polyesterification between citric acid and ethylene glycol. After drying at 110 °C for 12 h, the as-obtained powder was prefired at 500 °C for 1 h and then fully calcined at 950 °C for 7 h. The phase composition of the BSTF powders was examined by X-ray diffraction (Cu Kα, 45 kV, and 40 mA, PANalytical X’Pert Pro Multipurpose Diffractometer, Almelo, Netherlands) at room temperature in air. The BSTF powders exhibited a single phase cubic perovskite structure. Table 1 summarizes the compositions and notation of studied BSTF samples.

(3)

Intrinsic electronic defects, that is, electron−hole pairs, are generated via thermal ionization across the band gap:

nil ↔ e / + h•

[Oi//] − [V •• O] [STF]

where [STF] (or [BSTF]) is the number of STF (or BSTF) molecular formula units per unit volume.

where Oxo are oxygen ions on oxygen sites and Vxi are unoccupied oxygen interstitial sites in the form of uncharged interstitial vacancies with concentration equal to half the Fe content for the stoichiometric composition. The corresponding equilibrium constant is defined by K af/ =

(12)

(4)

with the equilibrium constant given by

K i = np

Table 1. Summary of BaxSr1−xTi1−yFeyO3−y/2+δ Sample Compositions and Notation

(5)

where n and p are concentration of electrons in the conduction band and holes in the valence band, respectively. The reduction reaction is expressed by / OOx ↔ V •• O + 2e +

1 O2 2

STF35 x = 0, y = 0.35

2 1/2 [V •• O ]n (pO2 ) [OOx ]

2 0 1/2 [V •• O ]0 n i (pO2 )

BSTF1005 x = 0.1, y = 0.05

BSTF5005 x = 0.5, y = 0.05

For TGA, the powders were pressed into pellets (weight of approximately 1 g) and sintered at 1200 °C to ∼70−80% relative density. Phase purity was confirmed by XRD and the respective XRD patterns are shown in Figure 1. The pellets were placed into an alumina crucible that was attached to a Pt wire hanging from one side of the microbalance beam (Cahn 2000). Controlled pO2 values were generated by use of O2−N2, H2−H2O−N2, and CO−CO2 gas mixtures, (using a total flow rate of 100 sccm over the sample with a

(7)

Alternatively, it is convenient to rewrite eq 7 in the form: 2 1/2 [V •• O ]n (pO2 )

BSTF5035 x = 0.5, y = 0.35

(6)

with the following mass action relation / K red =

BSTF1035 x = 0.1, y = 0.35

=1 (8)

with the denominator in eq 8 equivalent to [OxO]K/red (and 7 // x assuming [V•• O ]0 and [Oi ]0 are much less than [OO]). For this •• work, [VO ]0 and ni, correspond to the vacancy and electron concentration at the stoichiometric oxygen partial pressure pO02, that is, the pO2 where the concentrations of holes and electrons, and vacancies and interstitials, respectively, are equal (discussed later). As a result, from eqs 3 and 5, [V•• O ]0 and ni are defined as follows: 1/2 [V •• = (K af/ [OOx ][V ix])1/2 O ]0 = (K af )

n i = (K i)1/2

(9) (10)

In this case, the equilibrium constant for the reduction reaction (eqs 6 and 7), Kred, is ′ K red = (K af )1/2 K i(pO20 )1/2 = [OOx ]K red

Figure 1. XRD patterns of STF and BSTF pellets used for TGA measurements.

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Figure 2. Oxygen nonstoichiometry (plotted as δ) for (a) STF35, (b) BSTF1035, (c) BSTF5035, (d) BSTF1005, (e) BSTF5005 as a function of pO2 and temperature. The solid lines correspond to the fit based on the discussed defect model. constant 100 sccm N2 counter flow through the TGA electronics), monitored with a zirconia-based oxygen sensor. In an effort to alleviate the potential impact of carbonate formation in CO2 containing atmospheres, CO−CO2 measurements were performed last, followed by reoxidation to establish repeatability. Measurements using H2 were performed after O2−N2 conditions. The change in oxygen content Δδ with varying pO2 and temperature was calculated from the sample mass change, Δm, according to Δδ =

MSΔm MOm

reference measurements with an alumina sample of similar geometry. Based on the preferred oxidation state of Fe3+ and the resulting oxygen content of (3 − x/2 + δ), the measured plateau in oxygen nonstoichiometry was used as the reference point with δ0 = 0, so that the absolute value of δ was directly obtained from eq 13. The nonstoichiometry data were fitted based on the above equations with Kaf, Ki, and Kred as the fitting parameters. For a given oxygen vacancy concentration, n was calculated first by combining eqs 3, 11, and 12 and solving the quadratic equation for n with Kaf and Ki as the fitting parameters. pO2 could then be calculated based on eqs 7 and 8 with Kred as the fitting parameter. p, [O// i ], and δ were then determined from eqs 5, 3, and 13, respectively. Even though some acceptor-doped perovskite oxides show uptake of protons in humid atmospheres, significant dehydration occurs at temperatures above 700 °C,17 and this dehydration is more favorable for titanate materials (examined here) as compared to other common cerate and zirconate proton conductors.18,19 As the measured data for

(13)

with MS the molar mass of the sample, MO the molar mass of oxygen, and m the sample mass at room temperature in air. The error in mass measurement was ∼ ± 30 μg, which, for ∼1 g sample as used here, corresponds to an error in δ of ±0.0003. The buoyancy effect of gas and temperature on the measured mass change was corrected based on 2972



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our high temperature study (600−1000 °C) do not deviate from the expected defect equilibria behavior,7 no impact of proton solubility is expected in our material.

4. RESULTS Figure 2 shows the oxygen nonstoichiometry, δ, for STF35, BSTF1035, BSTF5035, BSTF1005, and BSTF5005 as a function of temperature and pO2. An oxygen-excess region (δ > 0) is observed at high pO2, followed by a plateau (δ = 0) at intermediate pO2 and an oxygen-deficient region (δ < 0) at lower pO2 for all compositions. The isothermal nonstoichiometry plots are shifted to lower pO2 with decreasing temperature, that is, reflecting a lower level of reduction at a given pO2 at reduced temperatures. Figure 3 compares the oxygen nonstoichiometry of BSTF at 800 °C with a fixed Fe content of 35 mol %, but varying Ba

Figure 4. Defect concentrations of STF35 and BSTF5035 as a function of pO2 at 800 °C.

The enthalpy of Frenkel defect formation is ∼0.5 eV independent of composition and in agreement with prior values determined from conductivity data for STF.7 The band gap energy, extracted from the defect model, are high temperature extrapolated values and are expected to be lower than values at 0 K7 and room temperature15 based on the band gap temperature dependence (i.e., decrease with increasing temperature).7 ΔHred for STF35 is also lower than the previously determined value (4.8 ± 0.1 eV),7 which, in the prior study, was based on a very limited number of nonstoichiometry data points. Finally, Figure 6 compares ΔHred and Eg for STF and BSTF with a constant iron content of 5 and 35 mol %, but with increasing Ba content. Upon addition of 10% Ba to STF, both energy terms decrease by ∼9−15% but do not change significantly further with additional Ba. Similar results are observed for BSTF with 5% Fe. These results are in agreement with our preliminary study.15 Figure 6 also shows that the difference in Eg values for 5 and 35 mol % containing BSTF is within the error bars.

Figure 3. Comparison of oxygen nonstoichiometry (plotted as δ) for STF35, BSTF1035, and BSTF5035 at 800 °C. The solid and dashed lines represent the fit of the defect model to the experimental data.

content. The nonstoichiometry isotherms are shifted to higher pO2 with increasing Ba content, shifting the onset of oxidation (high pO2 region) and reduction (low pO2 region) to higher pO2, which indicates enhanced reducibility with increasing substitution of Sr by Ba. In order to confirm the enhanced reducibility, ΔHred as well as Eg and the enthalpy for Frenkel pair formation, ΔHaf, were determined by applying the defect model discussed in the Theory section and developed in ref 7 to the experimental data. The solid and dashed lines in Figures 2 and 3 represent the fit of the defect model to the experimental data. Based on the thermodynamic relationships discussed in section 2, the concentrations of the defect species were calculated and are shown in Figure 4 as an example for STF35 and BSTF5035 at 800 °C. The stoichiometric oxygen partial pressure pO20 corresponds to the crossing point for n and p in Figure 4 and the inflection point in the nonstoichiometry plateau (Figures 2 and 3). ΔHred, Eg and ΔHaf were obtained from the slopes of Arrhenius plots of their respective equilibrium constants (Figure 5 a−c) and are summarized in Tables 2 and 3. Note that, aside from the changes in slope, the magnitudes of Kred and Ki follow the expected trends, that being an enhanced level of reduction and increased intrinsic electronic carrier density with increasing Fe (increased y from 0.05 to 0.35) and increasing Ba (increased x from 0 to 0.50). The latter case is particularly apparent for the 35 mol % Fe containing samples, where significantly larger nonstoichiometries aided in producing more reliable model fits to the data.

5. DISCUSSION The decrease in ΔHred and Eg confirm our hypothesis that the substitution of Sr in STF by the larger Ba should enhance reducibility and decrease the band gap energy. This is perhaps even clearer from the change in magnitudes of Kred and Ki with increasing x, as mentioned above. Through the enhanced concentration of oxygen vacancies and minority electron charge carriers (clearly shown in Figure 4 for increased Ba) for a given operating condition (T and pO2), an enhancement of D and kex, respectively, is expected. This should lead to improved cathode performance of BSTF and thereby make it an attractive candidate cathode material for intermediate-temperature SOFCs. However, other factors influencing cathode behavior, such as chemical stability of BSTF with respect to the solid electrolyte (e.g., YSZ) or to the gas phase (e.g., CO2, H2O), must be examined before its practical application can be assured.15 The present defect equilibria study allows one to estimate electron and hole mobilities from a prior electrical conductivity (σ) study on STF35,7 using the following equation. σi = qμi [i]

(14)

where q and μi are the elementary charge of an electron and mobility of charge carrier i. At 800 °C, this calculation yields an 2973



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also several orders of magnitude less than that reported for donor doped SrTiO3 (μe ∼ 6 cm2/(V s) via Ti 3d levels)20 at room temperature. The reduced electron mobility in the present material is a consequence of the electrons likely conducting predominantly via the much narrower Fe 3d levels, given that they lie lower than the Ti levels (see below) and are considerably diluted in concentration by Ti which makes up 65% of the B sites. This leads to small polaron type hopping between the localized Fe states with correspondingly low mobilities.21 The evidence of narrow Fe impurity bands is further supported by the low hole mobility relative to that of the related perovskite La1−xSrxFeO3−‑δ (μh ∼ 0.1−0.2 cm2/ (V s)) where the B-site is undiluted.22 Electronic conduction occurring largely via the Fe impurity levels is also supported by the small band gap energy derived in this study, as discussed next. An overestimated electron mobility led previously to an underestimation of the electron density, and in turn, an overestimation of the magnitude of the band gap energy (i.e., for STF35, Eg ∼ 1.38 eV in this study vs ∼2.6 eV in the previous study by Rothschild et al.)7 and is also likely attributed to the difference in band gap energy between the present work and that reported by Steinsvik and Norby, who followed a different defect modeling approach (∼2.2− 2.6 eV).10 The much smaller thermal band gap relative to that of SrTiO3 (3.17 eV)23 is believed to reflect the distance between the two Fe bands, derived from broadening of the Fe3+/4+ and Fe2+/3+ levels, lying within the original gap (which separates the top of the O2p derived valence band from the bottom of the Ti3d derived conduction band). The larger, optically derived band gap reported in ref 15 for BSTF remains consistent with optical measurements, where d-to-d transitions (between the Fe3+/4+ and Fe2+/3+) are forbidden (see Laporte selection rules).24 Future studies, probing the temperature and pO2 dependent electrical conductivity of BSTF are being pursued and should provide further insights regarding the impact that Ba substitution for Sr has on defect and transport properties in BSTF.

6. CONCLUSION The mixed conducting perovskite structured BSTF system, an extension of the STF model system, offers opportunities for composition and defect engineering of cathode materials of interest for intermediate-temperature SOFCs. Based on a defect chemical analysis of thermogravimetry derived nonstoichiometry data, the substitution of Ba for Sr on the A-site of STF was found to result in a decrease in both the reduction enthalpy and band gap energy. A consequent increase in minority carrier density and, ultimately, charge transfer rate is expected. The nonstoichiometry data obtained in this study also brought to light a previous underestimation of the minority electron density, by approximately a factor of 103. Our current efforts focus on investigating the cathode performance for STF and BSTF thin films to confirm the impact of adding Ba on k and D

Figure 5. Arrhenius plot of fitting parameters (a) Kred, (b) Ki, and (c) Kaf. The solid lines represent linear fits.

electron mobility of μe ≅ 0.0009 cm2/(V s) (at ∼10−21 bar O2) and a hole mobility of μh ≅ 0.005 cm2/(V s) (at ∼1 bar O2), the latter in agreement with previous estimates.7 μe, on the other hand, is ∼1000 times lower than a previous estimate, based only on electrical conductivity measurements7 and derived without the benefit of extensive nonstoichiometry data, as obtained for the first time in this study. This mobility is

Table 2. Enthalpy of Reduction, ΔHred, Band Gap Energy, Eg, and Enthalpy of Frenkel Defect Formation, ΔHaf, as a Function of Ba Doping for an Iron Content of 5 mol % ΔHred (eV) x

eV

0.1 0.5

3.55 ± 0.23 3.69 ± 0.10

ΔHaf (eV)

Eg (eV) −1

kJ mol

342.43 ± 21.99 355.74 ± 9.55

−1

eV

kJ mol

1.26 ± 0.12 1.10 ± 0.08 2974

121.48 ± 11.39 105.65 ± 7.62

eV

kJ mol−1

0.50 ± 0.001 0.56 ± 0.04

48.15 ± 0.1 53.94 ± 3.47



Article

Table 3. Enthalpy of Reduction, ΔHred, Band Gap Energy, Eg, and Enthalpy of Frenkel Defect Formation, ΔHaf, as a Function of Ba Doping for an Iron Content of 35 mol % ΔHred (eV) eV

kJ mol−1

eV

kJ mol−1

eV

kJ mol−1

0 0.1 0.5

3.893 ± 0.040 3.453 ± 0.107 3.521 ± 0.112

375.62 ± 3.86 333.16 ± 10.32 339.73 ± 10.81

1.379 ± 0.029 1.177 ± 0.029 1.204 ± 0.079

133.05 ± 2.79 113.56 ± 2.79 116.17 ± 7.62

0.518 ± 0.025 0.496 ± 0.065 0.564 ± 0.011

49.98 ± 2.41 47.86 ± 6.27 54.42 ± 1.06

(12) Jung, W.; Tuller, H. L. Solid State Ionics 2009, 180, 843−847. (13) Jung, W.; Tuller, H. L. ECS Trans. 2011, 35, 2129−2136. (14) Shannon, R. D. Acta Crystallogr. 1976, A32, 751−767. (15) Kim, J. J.; Kuhn, M.; Bishop, S. R.; Tuller, H. L. Solid State Ionics 2013, 230, 2−6. (16) Kroeger, F. A.; Vink, H. J., Solid State Physics, Seitz, F.; and Turnball, D., Eds., Academic Press, New York, 1956, 3, 307−435. (17) Kreuer, K. D. Solid State Ionics 1997, 97, 1−15. (18) Kreuer, K. D. Annu. Rev. Mater. Res. 2003, 33, 333−359. (19) Norby, T.; Larring, Y. Curr. Opin. Solid State Mater. Sci. 1997, 2, 593−599. (20) Baniecki, J. D.; Ishii, M.; Aso, H.; Kobayashi, K.; Kurihara, K.; Yamanaka, K.; Vailionis, A.; Schafranek, R. Appl. Phys. Lett. 2011, 99, 232111−3. (21) Tuller, H. L.; Nowick, A. S. J. Phys. Chem. Solids 1977, 38, 859− 867. (22) Søgaard, M.; Vang Hendriksen, P.; Mogensen, M. J. Solid State Chem. 2007, 180, 1489−1503. (23) Moos, R.; Hardtl, K. H. J. Am. Ceram. Soc. 1997, 80, 2549− 2562. (24) Atkins, P.; de Paula, J. Atkins’ Physical Chemistry, 7th ed.; Oxford University Press: Oxford, 2002; p 541.

Figure 6. Enthalpy of reduction and band gap energy as a function of Ba doping for an Fe content of 5 and 35 mol %. The dashed and dotted lines are a guide to the eye.

and to investigate the relative chemical stability of BSTF as a function of x and y.

■

ΔHaf (eV)

Eg (eV)

x

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This research was funded by the Department of Energy, Basic Energy Sciences under award DE SC0002633. J.J.K. thanks The Kwanjeong Educational Foundation for fellowship support. S.R.B. recognizes partial support from I2CNER, supported by the World Premier International Research Center Initiative (WPI), MEXT, Japan.

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REFERENCES

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Oxygen Nonstoichiometry and Defect Chemistry of Perovskite

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