A- and B-Site Ordering in the A-Cation-Deficient Perovskite Series La2

Article pubs.acs.org/cm

A- and B‑Site Ordering in the A‑Cation-Deficient Perovskite Series La2−xNiTiO6−δ (0 ≤ x < 0.20) and Evaluation as Potential Cathodes for Solid Oxide Fuel Cells Juan Carlos Pérez-Flores,*,† Domingo Pérez-Coll,‡ Susana García-Martín,§ Clemens Ritter,∇ Glenn C. Mather,‡ Jesús Canales-Vázquez,⊥ María Gálvez-Sánchez,⊥ Flaviano García-Alvarado,† and Ulises Amador*,† †

Facultad de Farmacia, Departamento de Química, Urbanización Montepríncipe, Universidad CEU San Pablo, Boadilla del Monte, E-28668 Madrid, Spain ‡ Campus Cantoblanco, Instituto de Cerámica y Vidrio-CSIC, C/Kelsen 5, E-28049 Madrid, Spain § Universidad Complutense de Madrid, Departamento de Química Inorgánica, E-28040 Madrid, Spain ∇ Institut Laue-Langevin, BP 156-38042, Grenoble Cedex 9, France ⊥ Renewable Energy Institute, University Castilla La Mancha, E-02006 Albacete, Spain S Supporting Information *

ABSTRACT: The La2−xNiTiO6−δ (0 ≤ x < 0.2) series has been investigated in order to assess its possible use as a solid oxide fuel cell (SOFC) cathode material. These perovskite-like oxides exhibit monoclinic symmetry, as determined by a series of high-resolution structural techniques (X-ray diffraction (XRD), neuron powder diffraction (NPD), selected-area electron diffraction (SAED), and transmission electron microscopy (TEM)). Ni and Ti order over the B-site and, unusually, for x > 0, the A-site ions are also ordered along the caxis in alternate La-rich and □-rich layers (where □ represents a vacancy). Structural determination combined with accurate compositional and magnetic characterization indicates a change in the predominant charge-compensating mechanism of A-site vacancies with composition. For x = 0.1, oxygen-vacancy formation seems to be the main-charge compensating mechanism, whereas, for x = 0.2, partial replacement of Ni by Ti in the Bsubstructure is dominant. In addition, a small amount of trivalent nickel is present in all samples. The composition dependence of the electrical conductivity of La2−xNiTiO6−δ (x = 0, 0.1, 0.2), investigated by impedance spectroscopy, as a function of temperature and oxygen partial pressure, is successfully interpreted on the basis of the relevant charge-compensating mechanisms and associated valence states. Thermal and chemical stability have also been studied in order to perform a preliminary electrochemical characterization as prospective cathode materials for SOFCs. The material La1.80NiTiO6‑δ exhibits excellent stability under oxidizing conditions and a polarization resistance of ∼0.5 Ω cm2 at 1073 K with a yttria-stabilized zirconia (YSZ) electrolyte, slightly lower than that of the state-of-the-art La1−xSrxMnO3 (LSM)-based cathodes. A higher thermal stability and a better chemical compatibility of La1.80NiTiO6−δ with common electrolytes (e.g., YSZ), in comparison with LSM, suggests that this oxide warrants further study and optimization as a prospective improved cathode material for SOFCs. KEYWORDS: SOFC, perovskite, cathode, LSM

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INTRODUCTION Perovskite-type oxides, with the general formula ABO3, are well-known materials with interesting applications as electrodes in several solid-state electrochemical devices for electrical power generation.1−4 Perovskites constitute a highly versatile and easily controllable family of compounds, from both the structural and chemical points of view.5 The structure accommodates a large variety of metal ions with variable oxidation state, and the anionic sublattice tolerates a considerable amount of vacancies. These characteristics render the properties of perovskite to be highly tunable in many cases. With respect to solid oxide fuel cell (SOFC) applications, © 2013 American Chemical Society

perovskites which exhibit both mixed ionic-electronic conductivity (MIEC) and high catalytic activity are of considerable interest as electrode materials.6−8 To date, perovskite oxides have been extensively used in the development of new electrode materials in both high- and intermediate-temperature solid-oxide fuel cells (IT-SOFCs). A simple strategy to improve the electronic conductivity is to induce mixed oxidation states in B-site ions, whereas oxide-ionic conduction can be enhanced Received: March 11, 2013 Revised: May 20, 2013 Published: May 21, 2013 2484

dx.doi.org/10.1021/cm4008014 | Chem. Mater. 2013, 25, 2484−2494

Chemistry of Materials

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ences in their structure, chemical reactivity, and electrical and electrochemical behavior are observed.

through the creation of anion vacancies. These effects can be achieved by appropriate aliovalent substitutions in the A- and/ or B-metal sites. In particular, in the most used cathode material, La1−xSrxMnO3 (LSM), the substitution of La3+ by Sr2+ improves electronic conductivity (due to small polaron hopping related to the mixed oxidation state of manganese, Mn3+/ Mn4+), with respect to the parent LaMnO3 phase. The range of oxygen nonstoichiometry, as a function of the oxygen partial pressure (pO2) is modulated, yielding a mechanically and chemically more stable material.9,10 However, some controversy concerning the optimal dopant level remains. Indeed, a compromise must be achieved between electrical conductivity and thermal expansion coefficient (and, therefore, the thermal stress in the cell), both of which are altered with the level of substitution.6,10,11 Nevertheless, despite the enormous possibilities that perovskite materials can offer, other factors must be taken into account, which limits their possible applications. These include the mechanical strength and chemical stability under reducing and/or oxidizing atmospheres, and undesirable side reactions at high temperatures with other components of the cell.12,13 For example, LSM reacts with yttria-stabilized zirconia (YSZ), which is the most commonly used electrolyte, above 1200 °C, limiting the sintering temperature that can be used for cosintering the electrolyte and the cathode.14−16 Although the diffusion of manganese into the electrolyte above this temperature can be considered negligible, the formation at the interface of the insulating phases La2Zr2O7, SrZrO3 and Sr2ZrO4,14 is an inherent problem of this cathode−electrolyte combination. These phases have even been detected at temperatures as low as 1000 °C.17−19 Among the huge variety of perovskites, LaNiO3-related compounds doped with various metals have attracted considerable attention as possible candidates for cathode materials, because of their high electronic conductivity.20,21 However, they exhibit low thermodynamic stability and decompose to La2NiO4 and NiO at high temperature and/or under reducing conditions.22,23 In addition, they present the compatibility problem of forming La2Zr2O7 with YSZ.17,24 Stability may be improved by partial substitution of Ni by other less-noble metals, such as Cr, Mn, Ti, and Ga;25 however, to date, mainly Co and Fe have been used for this purpose.26 Another approach for solving stability problems considers the creation of vacancies in the A-site.2,27−30 Indeed, some Adeficient perovskites exhibit high thermal stability and chemical compatibility while maintaining good electrochemical characteristics. Thus, La deficiency is a common practice for avoiding the formation of La2Zr2O7 by reaction of electrode materials with YSZ.31 In the present work, the La2−xNiTiO6−δ (0 ≤ x < 0.2) series has been investigated. A detailed chemical, structural, and microstructural characterization, using X-ray diffraction (XRD), neutron powder diffraction (NPD), selected-area electron diffraction (SAED), and high-resolution transmission electron microscopy (HRTEM) has been undertaken to reach a precise determination of its defect chemistry, in particular, the chargecompensating mechanisms of A-site deficiency. The electrical properties have been analyzed on the basis of the defect chemistry, and compared to those of the previously reported La2−xSrxNiTiO6−δ (0 ≤ x < 0.5) series.32 Although the two families of materials are apparently similar, important differ-

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EXPERIMENTAL SECTION

Sample Preparation. Synthesis of the La2−xNiTiO6−δ series (x = 0, 0.1, 0.2, and 0.33) was carried out using a modified Pechini method. Approximately 10 g of each composition were prepared by dissolving stoichiometric amounts of high-purity nickel (Aldrich, 99.99%) and La2O3 (Alfa Aesar, 99.9%) in ca. 20 mL of hot nitric acid (Panreac, 66%), before adding 50 mL of distilled water. Citric acid was subsequently added in a 3:1 molar ratio of citric acid:metal ions in conditions of constant heating and vigorous stirring. Insoluble TiO2 (anatase, Aldrich, 99.9%) was then added to obtain a homogeneous suspension. Once the volume had been reduced by half, 3 mL of diethylene glycol was added to promote polymerization. The resulting solid resin was cooled to room temperature, milled in an agate mortar and the powder burnt at 1073 K to remove the organic matter. After milling and homogenization, the resulting powder was fired at 1773 K for 24 h, then cooled to room temperature at a rate of ∼2 K min−1. Reduced samples were obtained on heating at 1225 K for 12 h (2 K min−1) under a dry H2(5%)/N2 stream, corresponding to a partial pressure of oxygen (pO2) of ∼10−21 atm. Characterization Techniques. Sample purity was determined by powder XRD (these data were also used for the structure characterization) on a Bruker D8 high-resolution diffractometer equipped with a position-sensitive detector (PSD) MBraun PSD-50 M using monochromatic Cu Kα1 (λ = 1.5406 Å) radiation obtained with a Ge primary monochromator. The angular range, step size, and counting times were selected to ensure the required data quality and resolution for structural refinement. The chemical composition of the samples was determined by energy-dispersive spectroscopy (EDS) using an EDAX detector on a FEI XL30 scanning microscope by analyzing ∼20 grains of every sample. For transmission electron microscopy (TEM), the samples were ground in n-butyl alcohol and ultrasonically dispersed. A few drops of the resulting suspension were deposited in a carbon-coated grid. Selected-area electron diffraction (SAED) studies were performed with a JEOL 2000FX electron microscope (double tilt (±45°)) operating at 200 kV. High-resolution transmission electron microscopy (HRTEM) studies were carried out with a JEM 3000F microscope operating at 300 kV (double tilt (±20°), point resolution of 1.7 Å), fitted with an EDS microanalysis system (Oxford INCA). Image calculations were performed using the Crystalkit and MacTempas programs based on the multislice approach. Neutron powder diffraction (NPD) experiments were performed at room temperature on the high-resolution D1A powder diffractometer at the Institut Laue-Langevin, (Grenoble, France). In order to access high Q-values, a monochromatic beam with a wavelength of 1.3893 Å was selected with a Ge monochromator, corresponding to an instrumental resolution within the range 2.7 × 10−3 ≤ (ΔQ/Q) ≤ 0.022. Structural refinements were carried out by the Rietveld method using the FullProf program33 by simultaneous fitting of XRD and NPD data. The neutron scattering amplitudes used in the refinement were 0.824, 1.03, −0.344, and 0.581 (× 10−12 cm) for La, Ni, Ti, and O, respectively; isotropic thermal factors (ITF) were used for all atoms. Constraints employed throughout refinement involved considering the perovskite B sites fully occupied and adopting the same thermal factor for all oxygen atoms. The oxidation state of nickel and oxygen content (assuming charge neutrality) of samples were determined by titration using potassium dichromate, as described in ref 34. Independent determinations were performed by thermogravimetric analyses using a D200 Cahn balance. Typically, ca. 70 mg of the sample were weighed to a precision of ±0.0005 mg at a total reduced pressure of 400 mbar containing 60% He and 40% H2. The sample was then heated to 1173 K at a rate of 5 K min−1. Total electrical conductivity was studied in several atmospheres on cylindrical pellets of 8 mm in diameter, 1−2 mm thick, with relative 2485



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densifications of >90% of the crystallographic density. Densified pellets were prepared by pressing the powders at 40 MPa and sintering at 1773 K for 6 h. Platinum electrodes were attached to both planar surfaces by painting with a platinum paste (Fuel Cell Materials) and annealing at 900 °C for 1 h to eliminate the organic content and ensure electrical contacts. Impedance spectroscopy was performed using an Autolab PGstat302N under continuous flows of 50 mL min−1 of air, O2, N2, and H2 (10%)/N2 (90%) in the range of 200−900 °C. Samples were equilibrated at 900 °C for 12 h in the corresponding atmosphere in order to reach equilibrium prior to the electrical study. Measurements were conducted at an amplitude voltage of 50 mV in the frequency range of 1−106 Hz. Isothermal measurements of electrical conductivity were also registered as a function of the oxygen partial pressure (pO2), as reported elsewhere.32 The system was first exposed to a continuous flow of 50 mL min−1 of H2 (10%)/N2 (90%) for 12 h at the corresponding temperature, in order to reach equilibrium under reducing conditions. The gas flow was then switched off and the system allowed to slowly recover the external pO2 over a period of several days. The electrical conductivity as a function of pO2 was evaluated during this time by impedance spectroscopy measurements registered in 10-min intervals. The chemical compatibility of the electrode materials based on La2−xNiTiO6−δ series with YSZ (Pi-KEM) and La0.8Sr0.2MnO3 (LSM, Next-Tech) was evaluated by mixing samples with YSZ or LSM in a 1:1 w:w ratio. Mixtures were pelletized and heated at 1373 K for 12 h, under both air and reducing conditions (H2 (5%)/Ar (95%)); phase analysis was then performed by XRD. Slurries for symmetrical-cell measurements were prepared by mixing the as-prepared materials and LSM powders with YSZ in a 1:1 w:w ratio with Decoflux (WB41, Zschimmer and Schwarz) as a binder. The electrodes were prepared by coating dense YSZ disks (thickness of ∼1 mm) with the slurries using active areas of 10 mm2. The resulting assemblies were fired at 1373 K for 10 h, then a platinum-based paste was deposited on top of the electrodes and the assemblies were refired at 1173 K for 2 h. Electrode microstructure was investigated using a JEOL scanning electron microscope (Model 6490L). Polarization measurements were performed with a two-electrode arrangement, as described previously,35 using flows of either pure oxygen or a H2 (5%)/N2 (95%) mixture in the temperature range of 973−1223 K. Impedance spectroscopy measurements of cells were carried out using a 1470E Solartron Cell Test in the frequency range of 10−106 Hz, using an amplitude voltage of 50 mV.

Table 1. Metal Contents of the La2−xNiTiO6−δ Samples Obtained by EDS-SEM and Compositions from NPD Element Content (at. %) nominal value of x

La

Ni

Ti

0.10

49.0(4)

24.1(2)

26.9(2)

0.20

48.2(7)

22.5(3)

29.3(4)

0.33

49.0(9)

20.4(3)

30.6(6)

experimental composition La1.92(3)Ni0.94(2)Ti1.06(2) O5.94 La1.86(3)Ni0.87(3)Ti1.13(3) O5.92 La1.92(4)Ni0.81(3)Ti1.19(3) O6.00

between the nominal and actual nickel and titanium compositions increases with increasing nominal A-site vacancy concentration. Therefore, it would seem that the structure cannot tolerate either a large amount of A-site or anionic vacancies. Hence, the concentration of La vacancies is significantly lower than the nominal value and excess Ti4+ replaces some Ni2+ to compensate the positive charge loss due to A-site vacancies. Since the actual composition of the La1.66NiTiO6−δ sample is very far from the nominal one, we will hereafter limit our study to samples in the range 0 ≤ x < 0.20. With regard to the earlier work, it should be stressed that the real limit of the solid solution series that we find is in good agreement with that proposed by these authors, with only slight differences, which can be attributed to the synthesis procedure used here. It is worth noting that the La2−xNiTiO6−δ series is stable in air up to the melting point, which is slightly above the synthesis temperature that we used, above which incongruent melting takes place with segregation of several phases. More important from a practical point of view is the stability under the reducing atmospheres used in SOFCs. On this point, our results diverge from those previously reported.30,36 In contrast to what is claimed by these authors, the La2−xNiTiO6−δ materials studied in this work are stable under reducing atmospheres. As an example, that of nominal composition La1.80NiTiO6−δ is stable after a 72-h treatment in an atmosphere of H2 (5%)/Ar (95%) (equivalent to a pO2 ≈ 10−21 atm) at 1323 K, as shown in Figure 1. Moreover, XRD patterns of the residues of the TGA experiments (not shown) in which a pO2 value as low as 10−38 atm at 1173 K was used revealed the samples to be stable, even under such extremely reducing conditions. Crystal Structure. SAED and HRTEM have been carried out with the objective of studying the existence of crystal defects or ordering effects in the title compounds, which are difficult to detect by means of XRD. The information provided by these techniques was used in the construction of a structural model for fitting the NPD and XRD data. Figure 2 shows SAED patterns along different zone axes of two crystals of nominal composition La1.80NiTiO6−δ. The patterns were indexed according to the ideal perovskite structure (cubic symmetry and lattice parameter, ap ≈ 4 Å). Tilting of the octahedral network to accommodate the size of the A-type cations is one of the most common distortions of the perovskite structure. Superlattice reflections in the SAED patterns associated to in-phase and/or antiphase tilting can help to determine the tilting system of the structure.37 In-phase tilt reflections clearly appear in the SAED patterns of La 1.80 NiTiO 6−δ . Antiphase-tilt superlattice reflections, 1 /2(ooo)p (where o denotes an odd Miller index) are also clearly observed. It is worth mentioning that rock-salt-type

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RESULTS AND DISCUSSION Compositional Range and Stability. Single-phase materials at the resolution of XRD were obtained for compositions 0 ≤ x ≤ 1/3 of the La2−xNiTiO6−δ series (see Figure SI 1 in the Supporting Information). This solid-solution range is greater than that reported in a previous paper30 in which the limit for La vacancies was established at x ≈ 0.14− 0.16. However, the study of our samples using NPD revealed the presence of some unreacted or segregated NiO, indicating off-stoichiometry in the analyzed samples. The second phase was also observed by SEM, the composition of which as confirmed by EDS. In addition, individual thin crystallites investigated by TEM (see below) were analyzed by EDS, confirming that the actual stoichiometries of crystallites were somewhat different from the nominal values. The actual compositions of bulk materials, listed in Table 1, correspond to EDS results obtained via SEM. It can be seen that the Ti/Ni ratio is always greater than the nominal one (1:1). Table 1 also shows the NPD refined composition, which confirms that the actual stoichiometries exhibit a Ti excess and a Ni deficiency for the La-deficient samples. It can also be seen that the difference 2486



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Figure 1. Experimental (red circles) and calculated (black continuous line) XRD patterns (and their difference, blue line at the bottom) for La1.80NiTiO6−δ after a 72-h heat treatment in an atmosphere of H2(5%)/Ar(95%) at 1323 K, assuming the same structure as the as-prepared material (see Table SI 1 in the Supporting Information). Green vertical bars indicate the positions of the Bragg peaks of the phases contained in the sample.

Figure 2. Selected-area electron diffraction (SAED) patterns of a La1.80NiTiO6‑δ crystal along (a) the [001]p zone axis and (b) the [11̅ 0]p zone axis. (c) SAED pattern of a thin crystal of La1.80NiTiO6−δ along the [11̅ 0]p zone axis.

Figure 3. (a) High-resolution transmission electron microscopy (HRTEM) images of a La1.80NiTiO6−δ crystal along the [1̅10]p zone axis and the corresponding fast Fourier transform (FFT); (b) calculated HRTEM image inserted in the experimental one.

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Figure 4. Experimental (red circles) and calculated (black continuous line) NPD patterns (and their difference, blue line at the bottom) for La1.80NiTiO6−δ; the structure model is given in Table SI 1 in the Supporting Information. Green vertical bars indicate the positions of the Bragg peaks of the phases contained in the sample; the second row corresponds to a small amount of NiO.

0 0) and 2c (1/2 0 1/2), which allows an ordered arrangement of the B-cations. Since the neutron scattering amplitudes of Ni and Ti are significantly different, the site occupancies can be readily identified; this is not the case for XRD, since the X-ray scattering factors of Ti and Ni are too similar. In the parent compound La2NiTiO6, Ni2+ and Ti4+ are ordered, although some degree of intermixing or antisite defects (AS) of Ni and Ti on the B′ and B″ sites, respectively, is present (∼8%). This sample was found to be slightly Ni-deficient and a small amount of NiO (∼1 wt %) was detected by NPD (but not by XRD). This Ti overstoichiometry, accompanied by a small amount of segregated NiO, is also found in all samples of the La2−x NiTiO6−δ series, and seems to be very common in Adeficient nickelates.2,27−30 Accordingly, we used the compositions shown in Table 1 for fitting the diffraction data. Table SI 1 in the Supporting Information collects the refined structural parameters of the La2−xNiTiO6−δ series (x = 0.1 and 0.2), as obtained from XRD and NPD data. Selected structural information is given in Table SI 2 in the Supporting Information. All members of the La2−xNiTiO6−δ series exhibit perovskite tolerance factors (t) below unity (see Table SI 2 in the Supporting Information), leading to a tilting of the B′O6 and B″O6 octahedra. According to the Glazer notation,40 the tilting system for these compounds is (a−a−c+), indicating that the octahedra rotate along the Cartesian axes x and y in consecutive layers in opposite directions, whereas along the z-axis they rotate in the same direction. The octahedral tilts can be calculated from the B′−O−B″ bond angles (ω) as (180 − ω)/ 2; the values obtained from NPD are given in Table SI 2 in the Supporting Information. In the La2−xNiTiO6−δ series, the perovskite tolerance factor slightly decreases with x and, as a result, the structure tends to be more distorted, to accommodate a significant degree of A-ion vacancies (see Table SI 2 in the Supporting Information); this effect is particularly evident in the octahedral tilting angles.

ordering of the Ni and Ti could give rise to the strong 1 /2(111)p reflection, indicated with an arrow in the patterns of the [1̅10]p zone axis (Figure 2b). The intensity of this reflection suggests a combination of tilting of the octahedra and Ni and Ti ordering. Rock-salt-type ordering of Ni and Ti was previously found in La2NiTiO6 by neutron diffraction; however, the ordering is suppressed when 5% of La is substituted by Sr in La2−xSrxNiTiO6−‑δ.32 Most likely, ordering of Ni and Ti is exhibited in the title compounds and is associated with ordering in the oxygen substructure and/or on the perovskite A-site. The combined use of different diffraction techniques is needed to confirm this hypothesis. Some reflections, such as 1/2(00l)p, probably arise as a consequence of dynamical effects, because they do not appear in the patterns of thin crystals (Figure 2c). Figure 3a shows the HRTEM image of the [1̅10]p zone axis and the corresponding fast Fourier transform (FFT). Contrast differences corresponding to a periodicity of 2ap × √2ap are observed. This indicates that two different positions should be considered for the B atoms, because of cation ordering or displacements from the center of the octahedra. In TEM images of perovskites with a √2ap × √2ap × 2ap superstructure due only to octahedral tilting, the √2ap periodicity is not observed, because the effects associated with the anion substructure generally do not give strong contrast differences.38 As mentioned above, in a previous paper,32 we confirmed the structure of La2NiTiO6 to be monoclinic (space group P21/n, #14) from NPD data, with a unit cell of dimensions √2ap × √2ap × 2ap and not orthorhombic (space group Pnma, #62) as previously reported.39 The reason for this reduction in symmetry is mainly due to the ordering of Ti and Ni in the B′ and B″ positions of the structure, although the monoclinic angle is very close to 90°. In the orthorhombic model, there is only one crystallographic site, 4b (0 0 1/2), for B cations, i.e., the symmetry does not permit an ordered B-cation arrangement. In contrast, in the monoclinic model, the 4b site in Pnma can be split into two independent crystallographic sites, 2d (1/2 2488



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□-rich layers to compensate for the loss of positive charge due to A-site vacancies. The proposed structural model with B-cation rock-salt-type ordering, layered ordering of vacancies in the A positions, and octahedral tilting (Table SI 1 in the Supporting Information) was adopted to perform TEM image simulations. Figure 3b depicts the calculated image for a 400 Å specimen thickness and a −300 Å defocus of the objective lens, inserted into an experimental image of the [11̅ 0]p zone axis for comparison. Good agreement between the simulated and experimental images is observed. The image simulation seems to confirm that the contrast difference observed in the HRTEM images, indicating that √2ap periodicity is mainly associated to the Ni/ Ti rock-salt-type ordering. The existence of A-site vacancies is relevant to understanding the electrical properties of the La2−xNiTiO6−δ series, since they are charge-compensated by both the formation of vacancies in the anion substructure (given by eq 1 in Krö ger-Vink notation):

Interestingly, the change from monoclinic to orthorhombic symmetry observed upon substitution of La3+ by Sr2+ is not found upon A-site vacancy creation. As previously discussed, this reduction in symmetry is only related to the 3d metals ordering. Ni and Ti remain ordered, and the entire La2−xNiTiO6−δ series displays a monoclinic symmetry with a β angle of ∼90° (see Table SI 1 in the Supporting Information). Ordering is made evident by the presence of the (011), (101) and (10̅ 1) peaks located at 17.6° 2θ in the NPD pattern shown in Figure 4 (these indices are given in the so-called “diagonal perovskite cell” √2ap × √2ap × 2ap and correspond to (1/2 1/2 1/2)p, (1/2 −1/2 1/2)p and (−1/2 1/2 1/2)p, respectively, in the cubic perovskite cell). The intensities of these peaks increase as the difference in the scattering powers between the B′ and B″ cations increase. The high definition of these peaks indicates that the Ti/Ni ordering is long-range. Moreover, intermixing of Ni and Ti over the B-sites is observed to some extent for the entire series, to a similar degree as that found in the parent compound La2NiTiO6. Metal−oxygen distances and bond valence sums (BVSs) (see Table SI 2 in the Supporting Information) support the cationic distribution determined by NPD (Table SI 1 in the Supporting Information). Indeed, for La2NiTiO6, most of the smaller Ti4+ cations are located on the B′-sites, giving short B′−O distances and high BVSs in this position. Conversely, Ni2+ is located on the B″-sites, as revealed by longer B″−O distances and lower BVSs. As the amount of both La vacancies and titanium increases, Ti4+ and Ni2+ remain mostly ordered, and the B′−O and B″−O distances and corresponding BVSs remain clearly different. As stated above, the creation of La vacancies does not induce disordering of the 3d metals, whereas, in the La2−xSrxNiTiO6−δ series, Ni and Ti are disordered for a degree of substitution as low as x ≈ 0.15. In the latter, a random distribution of Sr and La (and possibly of the oxygen vacancies created by aliovalent substitution) will provide the driving force for disordering of Ni and Ti. In this regard, it is well-established that, in A-site-vacant perovskites closely related to the title oxides, such as La0.6Sr0.1TiO3,41 the A-cations (La3+ and Sr2+) and vacancies order in layers, which are alternately vacancy-rich and vacancypoor. Furthermore, many other A-site-vacant perovskites, such as RE1−xMO3 (RE = La, Ce, Pr, Nd; M = Ti (x = 1/3), Nb (x = 2 /3)),42,43 and La(2/3)−xLi3xTiO344 exhibit a layered order of A ions and vacancies. B-cation order is quite usual for A2BB′O6 stoichiometries.45 As stated above, ordering of A ions is also frequent in A1−x□xBO3 A-site deficient perovskites and some perovskites with two types of A and B ions, denoted as AA′BB′O6, are known to present order in both the A and B substructures.46,47 The common features of the three types of structure are rocksalt ordering of the B-cations, layered ordering of the A-site cations, and distortions involving octahedral tilting. Some or all of these characteristics can be present in a given compound. We thus refined a structural model in which all these structural features coexist, together with the appropriate concentration of oxygen vacancies; the final structural parameters for this model are given in Table SI 1 in the Supporting Information. The refinement reveals that, besides rock-salt ordering of Bsite ions (Ti and Ni) suggested by SAED and HRTEM, the Asite ions are also ordered in alternate La-rich and □-rich (where □ denotes a vacancy) layers perpendicular to the 2ap cell parameter. The B-cations are displaced along the c-axis from the equatorial planes of the BO6 octahedra toward the

La ×La +

3 × 3 1 OO → V‴La + V •• La 2O3 O + 2 2 2 2+

and oxidation of Ni

2La ×La + 6Ni ×Ni +

to Ni

(1)

3+

3 O2 → 2V‴La + 6Ni•• Ni + La 2O3 2

(2)

The presence of oxygen vacancies in La2−xNiTiO6−δ has been confirmed by NPD. Table SI 1 in the Supporting Information shows that the oxygen content decreases as the La-vacancy content increases. Although the refined oxygen contents and those calculated from the refined cationic compositions (assuming charge neutrality and La, Sr and Ti in their most stable oxidation states (3, 2 and 4, respectively) and Ni in its actual oxidation state determined by TGA and chemical titration, as discussed below) do not fully agree in all cases, both series of values display the same trend. Nickel Oxidation State. Redox titration and thermogravimetric analyses suggest that the amount of Ni 3+ in La2−xNiTiO6−δ is low and similar for the entire series. The percentage of trivalent nickel as determined by chemical titration is ca. 0.7(3) for all samples, in agreement with the result given by TGA. Therefore, oxygen-vacancy formation seems to be the dominant charge-compensating mechanism throughout the La2−xNiTiO6−δ series. Interestingly, in the related La2−xSrxNiTiO6−δ series,32 oxygen-vacancy creation also dominates, but the amount of Ni3+ is higher (ca. 0.7% to 3%). Magnetic Properties. The susceptibility versus temperature (χ vs T) data for the La2−xNiTiO6−δ series with x = 0, 0.1, and 0.2 are presented in Figure 5. All compounds are paramagnetic above 75 K and obey the Curie−Weiss law. Well-defined maxima appear at ∼25 K for the series (a slight decrease in temperature is observed as x increases) suggesting long-range antiferromagnetic (AF) ordering, in agreement with previous results.32,48 The presence of La vacancies has little or no detrimental effect on the magnetic ordering; a magnified view of the peaks in the χ vs T plot is presented in the inset of Figure 5. Since, in this type of material, the magnetic order is strongly hindered by B-site disorder, these results confirm the long-range ordering of Ni and Ti within the B′ and B″ sites, as determined by SAED and NPD (see Table SI 1 in the Supporting Information). Interestingly, again, the A-site-vacant perovskites derived from La2NiTiO6 behave differently to the Sr-substituted series La2−xSrxNiTiO6−δ. In the latter family, 2489



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vacancies are created, the magnetic moment remains above the SO value but slightly decreases along the series with increasing x. This suggests that (a) the oxidation state of Ni ions is ∼2, (b) the quenching of the orbital moment is not complete, and (c) a small fraction of nickel is in the trivalent state. Quenching of the orbital moment seems to be related to the degree of order in the B-sites. Hence, in La2−xSrxNiTiO6−δ, Ni and Ti are disordered over the B-sites and SO moments are observed, whereas, in La2−xNiTiO6−δ, 3d metal order is preserved and magnetic moments well above the SO value are obtained. The slight decrease of the overall effective moments with x supports the presence of small amounts of trivalent nickel, as indicated by NPD, TGA, redox titration, and electrical measurements (see the next section). Electrical Properties. Figure 7 shows the Arrhenius representation of total conductivity (σT) in air for different Figure 5. Susceptibility (χ) versus temperature (T) for materials of the La2−xNiTiO6−δ series.

magnetic order associated to Ni/Ti ordering becomes shortrange for doping levels as low as x = 0.1 (and completely disappears for x ≈ 0.15), whereas the presence of La vacancies preserves both the 3d metals order and the AF coupling for the entire series. This is most likely due to the ordered arrangement of La and vacancies in layers along the c-axis, as indicated by the NPD data. The overall effective moments for the La2−xNiTiO6−δ series obtained by Curie−Weiss fits (χ(T) = C/(T − θ)) are plotted in Figure 6; the calculated moments as a function of x,

Figure 7. Arrhenius representation of the total conductivity in air for the La2−xNiTiO6−δ system.

compositions in the La2−xNiTiO6−‑δ system (x = 0, 0.1, 0.2). All samples exhibit an almost-linear behavior in the studied temperature range, suggesting that the oxygen nonstoichiometry is mainly independent of temperature, as previously found for La0.95Ni0.5Ti0.5O3−δ.36 The high values of activation energies, in the range of 0.78−0.89 eV, are consistent with that reported for the parent phase, La2NiTiO6,32 and are very close to the value of 0.92 eV found for La0.95Ni0.5Ti0.5O3−δ.30 The high activation energy may be an indication of strong trapping of electronic defects. The introduction of vacancies in the La position increases the conductivity in air of composition x = 0.1, compared to that of x = 0. The conductivity of x = 0.1 obtained in O2, air, and N2 (Figure 8) shows that the total conductivity increases as the oxidizing character of the gas becomes greater. These results indicate that transport is governed by p-type carriers (electron holes) and increases as the oxygen partial pressure (pO2) increases, according to 1 × • O2 + V •• O → OO + 2h (3) 2 Although the main charge-compensation process associated with the creation of La vacancies in the crystal structure is suggested to be the generation of oxygen vacancies, the

Figure 6. Effective magnetic moments for La2−xNiTiO6−δ (0 ≤ x < 0.2) and those calculated assuming fully oxidation of Ni2+ upon creation of La vacancies.

assuming oxidation of Ni2+ as a charge-compensation mechanism (eq 2), is also plotted in this figure. The Weiss constant is approximately −140 K in all cases, suggesting strong AF interactions. The paramagnetic moments displayed in Figure 6 deserve some comments. For undoped La2NiTiO6, the experimental value of 3.09(2) μB is in good agreement with previous reports32,48 and with that expected for the octahedral field term 3A2 of Ni2+ (3.2 μB).49 However, this value is far above the spin-only magnetic moment for a d8 cation such a Ni2+ (2.83 μB); an incomplete quenching of the orbital moment is very often observed for these ions.32 When La 2490



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Figure 9. Total conductivity in H2(10%)/N2(90%) of samples with x = 0, 0.1, and 0.2.

Figure 8. Arrhenius plot of total conductivity of La1.9NiTiO6−δ under a continuous flow rate of 50 mL/min of O2, air, and N2.

existence of a minor Ni3+ content may be responsible for the ptype electronic conductivity (eq 2). However, further increasing the content of La vacancies to x = 0.2 has the effect of decreasing the conductivity to values that are in the range of the La-stoichiometric sample (Figure 7). These results are inconsistent with the simple model represented by eq 2, in which a higher content of La vacancies is expected to produce an increase of the p-type-carrier concentration ([Ni3+]). Several factors could be responsible for this unexpected behavior. As mentioned previously, there is a slight Ti overstoichiometry, which produces segregation of a small amount of NiO (see Table SI 1 in the Supporting Information). This Ti overstoichiometry may compensate the creation of La vacancies, according to 3x × 3x •• Ni Ni → x V‴La + TiNi x La ×La + (4) 2 2 2+ 4+ × •• where NiNi represents Ni and TiNi represents Ti , both in the B″-position of the perovskite predominantly occupied by divalent nickel (see Table SI 1 in the Supporting Information). This charge-compensation mechanism, as opposed to that represented in eq 2, leads to a lower concentration of p-type carriers (present as Ni3+ species) for higher Ti overstoichiometry, as found for x = 0.2 in comparison to x = 0.1 (see Table 1), which decreases the electronic conductivity of the former (Figure 7). O n th e o th er h an d, at om istic sim ula tio ns o f La1−xNi0.5Ti0.5O3−δ suggested a strong interaction between oxygen vacancies, lanthanum vacancies, and nickel(III) cations, yielding ternary clusters, which prevent the mobility of charge • carriers such as V•• O and Ni , and may be responsible for the low conductivity of these compounds,30 which is predominantly dominated by p-type carriers.36 As a consequence, a higher La deficiency may produce a reduction in the concentration of untrapped, mobile species, thus decreasing the total conductivity and increasing the activation energy (Figure 7). Under reducing conditions, the electrical behavior is modified, as is apparent in Figure 9, which shows the Arrhenius representation of conductivity under a continuous flow of 50 mL min−1 of H2(10%)/N2(90%). Although the samples present

almost-linear behavior, with activation energies in the range of 0.59−0.69 eV, the introduction of La deficiencies in the La2−xNiTiO6 system produces a decrease in the total conductivity, as is observed for x = 0.1; however, a higher content of La vacancies (x = 0.2) improves the total conductivity, with respect to x = 0.1. Therefore, the oxygen loss can be expressed as follows: 1 × OO → O2 + V •• O + 2e′ (5) 2 where the electrons are located at titanium positions as Ti3+, according to 2Ti ×Ti + 2e′ → 2Ti′Ti

(6)

• 2Ti•• Ni + 2e′ → 2TiNi

(7) 4+

Two reduction processes could take place for Ti ions due to Ti overstoichiometry and the ordering of B cations, given by eqs 6 and 7. A simple analysis of eqs 5−7 indicates that the total conductivity under reducing conditions should be mainly governed by n-type charge carriers. This is consistent with the monotonous increase in conductivity with decreasing oxygen partial pressure in the low-pO2 range (Figure 10). The determination of the pO2 dependence of conductivity was performed using a dynamic technique which commonly presents problems of interpretation in the intermediate range of pO2 in which experimental data correspond with nonequilibrium conditions.50 This inconvenience prevents the identification of the n−p transition in the σT−pO2 diagram, and only the more-extreme conditions should be taken into consideration. The mass action constant of eq 5 can be expressed as follows: 2 K O = pO 1/2 [V •• O ]n 2

(8)

where n corresponds to the concentration of electrons residing as Ti3+ cations. Assuming that the main compensation mechanism for the creation of La deficiency is the 2491



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ibility of the sample with x = 0.2, because of the lower concentration of free oxygen vacancies.30 Electrochemical Performance. Chemical Stability and Compatibility. The chemical compatibility of the La2−xNiTiO6−δ series with YSZ and LSM used as electrode components was evaluated. The title compounds react with YSZ under reducing conditions, as shown in Figure SI 2 in the Supporting Information for x = 0.20. This limits the use of these materials in the anode side of a SOFC. In contrast, when La2−xNiTiO6−δ/YSZ and La2−xNiTiO6−δ/ LSM (1:1) mixtures are heated in air at 1373 K simulating cathodic conditions, no decomposition or reaction occurs (see Figures SI 3a and SI 3b in the Supporting Information for La1.80NiTiO6−δ). Thus, the possible use of the title materials in the cathode side should present no chemical compatibility problems. Polarization Tests on Symmetrical Cells. Since the La2−xNiTiO6−δ/YSZ (1:1) composites are not stable in reducing atmospheres, the results of polarization measurements of symmetrical cells under flowing 5%H2/Ar are not discussed in detail. We note, however, that their polarization resistances are rather large, compared to Ni/YSZ-based anodes (7.6 and 30.5 Ω cm2 at 1073 K for x = 0.10 and x = 0.20), which, together with their poor chemical stability, precludes their use as fuel electrodes (anode) in SOFCs. The thermal evolution of the polarization resistance of a symmetrical cell consisting of composites La2−xNiTiO6−δ/YSZ electrodes over YSZ electrolyte in pure oxygen is represented in Figure 11a. The polarization resistance exhibits an Arrheniuslike temperature dependence, with an activation energy of Ea ≈ 1.33 eV. It can be observed that the polarization resistance decreases with an increasing number of A-site vacancies, with the lowest Rp value being obtained for x = 0.20. The La1.80NiTiO6−δ-based electrodes exhibit polarization resistances of ∼0.5 Ω cm2 at 1073 K, which is a slightly better value than that of the state-of-the-art LSM-based cathodes (also shown in Figure 11a) and in a range approximating the demands of SOFC manufacturers, i.e., 0.1−0.15 Ω cm2. On the other hand, the addition of LSM to the La2−xNiTiO6−δ/YSZ electrodes (1:1:2 weight ratio) with the aim of increasing the electrical conductivity of the resulting composite, Figure 11b, has a minor effect on the electrode-polarization resistance. We note that the LSM powder used as a reference is a commercial product whose synthesis conditions are fully optimized to reach a narrow particle size distribution (0.3− 0.5 μm, large surface area (6 m2/g), etc), whereas the title material is reported for the first time and obviously must be optimized in the same way. Moreover, LSM reacts with YSZ at high temperatures (ca. at 1373 K), 14,16,17 whereas La1.80NiTiO6−δ/YSZ composites are stable at these temperatures. In addition, the processing of LSM/YSZ composites and cell assemblies can be considered as standard procedures; we used the same electrode composition and methods for our material, but they are not necessarily the optimum conditions in this case. The performance of the La1.80NiTiO6−δ/YSZ composites may be greatly enhanced, therefore, by optimizing the processing conditions, and improving that of LSM/YSZ composites. A cross-sectional view of the La1.80NiTiO6−δ-based cell after operation is presented in Figure 12. A good adherence between the electrolyte and the electrode is observed with no cracks or delamination occurring at the interfaces. Moreover, the electrode exhibits a porous microstructure which facilitates

Figure 10. Conductivity (σT) as a function of the oxygen partial pressure (pO2) for x = 0.1 and 0.2.

incorporation of oxygen vacancies (eq 1), the electroneutrality condition is expressed as 2[V •• O ] ≈ 3[V‴ La] + n

(9)

The usual observance of a −1/4-power-law dependency for conductivity versus pO2 only holds under the assumption that oxygen vacancies are constant and fixed by the content of La deficiency, according to 2[V •• O ] ≈ 3[V‴ La]

(10)

This relation dictates that the concentration of electronic carriers is expressed as 2K O p −1/4 n≈ 3[V‴La] O2 (11) On the other hand, if the oxygen-vacancy concentration is mainly governed by the oxygen equilibrium (eq 5), then 2[V •• O] ≈ n

(12)

The concentration of electron species can then be represented as follows:

n ≈ 2K OpO −1/6 2

(13)

For a general situation in which neither of these scenarios predominate, the slope lies between −1/4 and −1/6. In Figure 10, the low-pO2 behavior shows a −1/6-power-law dependence, whereas for the intermediate range of pO2 (∼10−12 atm), a transition to a −1/4 dependence is apparent. This latter behavior is attributable to a dominating mechanism in which the creation of La vacancies is charge-compensated by the creation of oxygen vacancies. Accordingly, this process should inhibit the reduction of Ti4+ to Ti3+, as is inferred from eqs 5−7], decreasing the n-type electronic conductivity, compared to the La stoichiometric sample (Figure 9). However, the slightly higher Ti overstoichiometry in the sample should lead to a greater reduction of Ti4+ to Ti3+ under reducing conditions, as is evident from combining eqs 4 and 7, thus increasing the ntype conductivity, compared to that of x = 0.1 (see Figures 9 and 10). Note that the higher concentration of the aforementioned ternary clusters should also improve reduc2492



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may help to further improve the microstructure and lower the polarization resistance. In a previous paper, we reported that La2−xSrxNiTiO6−δbased electrodes51 exhibit polarization resistances of >1.5 Ω cm2 under oxidizing conditions at 1073 K. In comparison to strontium doping, the creation of vacancies in the perovskite Asite seems to be a more favorable strategy for improving the electrochemical performance of this system. Further improvements that could be achieved through optimized electrode fabrication and morphology would render the title perovskites useful cathode materials for SOFCs.

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CONCLUSIONS The La2−xNiTiO6−δ (0 ≤ x < 0.2) series has been investigated as prospective cathode material for SOFCs. Crystal structure was determined by high-resolution structural techniques (XRD, NPD, SAED, and TEM). The entire series exhibits monoclinic symmetry; for the parent compound, the ordering of the B-ions makes the space group P21/n (#14), whereas, for the other members of the series, both A- and B-site ordering occurs, reducing the symmetry to P21 (#4). Ni and Ti order over the Bsites in a rock-salt-like arrangement and the A-site ions order along the c-axis in alternating La-rich and □-rich layers (where □ is a vacancy). Structural determination combined with accurate compositional and magnetic characterizations indicates a change in the charge-compensating mechanism of A-site vacancies as a function of x. For x = 0.1, oxygen vacancy formation is prevalent whereas, for x = 0.2, partial replacement of Ni by Ti in the B-substructure is dominant. The electrical properties of the compounds La2−xNiTiO6−δ (x = 0, 0.1, 0.2), investigated by impedance spectroscopy as a function of temperature and oxygen partial pressure, reveal an inconsistent variation of electrical conductivity with x. This behavior can be rationalized on the basis of the prevailing defect chemistry, involving oxygen vacancies, and small amounts of Ti-overstoichiometry and Ni3+. The thermal and chemical stability of the title compounds have been also studied in order to perform a preliminary electrochemical characterization as prospective cathode materials for SOFCs. Composites of La1.80NiTiO6−δ/YSZ have excellent stability under oxidizing conditions and exhibit polarization-resistance values of ∼0.5 Ω cm2 at 1073 K, which is slightly better than the state-of-the-art LSM-based cathodes. The high thermal stability and better chemical compatibility of La1.80NiTiO6−δ with YSZ (commonly used as SOFC electrolyte) in comparison to commercial LSM, (and similar to A-site-deficient LSM31), suggests that this oxide warrants further study as a prospective improved cathode material for SOFCs. In particular, more competitive performances as SOFC cathode may be achieved through optimization of the electrode processing.

Figure 11. Temperature dependence of the polarization resistance of symmetric cells based on the La2−xNiTiO6−δ system on YSZ electrolyte under cathodic conditions: (a) electrodes made by combining the material with YSZ and (b) electrodes made by combining with YSZ+LSM.

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ASSOCIATED CONTENT

S Supporting Information *

Table SI 1: Refined structural parameters for La2−xNiTiO6−δ (0 ≤ x < 0.20) materials as obtained from XRD and NPD. Table SI 2: Selected structural information for La2−xNiTiO6−δ (0 ≤ x < 0.20). Figure SI 1: XRD patterns of La1.80NiTiO6−δ at 1773 K for 24 h. Figure SI 2: XRD pattern of a 1:1 w:w mixture of La1.80NiTiO6−δ and YSZ heated under reducing conditions at 1373 K for 12 h. Figures SI 3a and SI 3b: XRD patterns of 1:1 w:w mixtures La1.80NiTiO6−δ/YSZ and La1.80NiTiO6‑δ/LSM,

Figure 12. SEM micrograph of the cathode La1.80NiTiO6−δ/YSZ prepared at 1373 K overnight.

transport of the gas species to and from the active sites at the triple phase boundaries. Nevertheless, the use of pore formers 2493



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(19) Tricker, D. M.; Stobbs, W. M. In Proceedings of the 14th Risø International Symposium on Materials Science: High Temperature Electrochemical Behaviour of Fast Ion and Mixed Conductors; Poulsen, F. W., Bentzen, J. J., Jacobsen, T., Skou, E., Ostergard, M. J. L., Eds.; Risø National Laboratory: Roskilde, Denmark, 1993; p 453. (20) Hrovat, M.; Katsarakis, N.; Reichmann, K.; Bernik, S.; Kus̆cĕ r, D.; Holc, J. Solid State Ionics 1996, 83, 99. (21) Skinner, S. J. Int. J. Inorg. Mater. 2001, 3, 113. (22) Drennan, J.; Tavares, C. P.; Steele, B. C. H. Mater. Res. Bull. 1982, 17, 621. (23) Zinkevich, M.; Aldinger, F. J. Alloys Compd. 2004, 375, 147. (24) Kostogloudis, G. C.; Tsiniarakis, G.; Ftikos, C. Solid State Ionics 2000, 135, 529. (25) Chiba, R.; Yoshimura, F.; Sakurai, Y. Solid State Ionics 1999, 124, 281. (26) Echigoya, J.; Hiratsuka, S.; Suto, H. Mater. Trans., JIM 1989, 30, 10. (27) Konysheva, E.; Irvine, J. T. S. Chem. Mater. 2011, 23, 1841. (28) Konysheva, E. Y.; Xu, X.; Irvine, J. T. S. Adv. Mater. 2012, 24, 528. (29) Knudsen, J.; Friehling, P. B.; Bonanos, N. Solid State Ionics 2005, 176, 1563. (30) Yakovlev, S. O.; Kharton, V. V.; Naumovich, E. N.; Zekonyte, J.; Zaporojtchenko, V.; Kovalevsky, A. V.; Yaremchenko, A. A.; Frade, J. R. Solid State Sci. 2006, 8, 1302. (31) Minh, N. Q.; Takahashi, T., Eds.; Elsevier: Amsterdam, 1995. (32) Perez Flores, J. C.; Ritter, C.; Perez-Coll, D.; Mather, G. C.; Garcia-Alvarado, F.; Amador, U. J. Mater. Chem. 2011, 21, 13195. (33) Rodríguez-Carvajal, J. Physica B (Amsterdam, Neth.) 1993, 192, 55. (34) Yuste, M.; Perez-Flores, J. C.; de Paz, J. R.; Azcondo, M. T.; Garcia-Alvarado, F.; Amador, U. Dalton Trans. 2011, 40, 7908. (35) Ruiz-Morales, J. C.; Canales-Vázquez, J.; Peña-Martínez, J.; López, D. M.; Núñez, P. Electrochim. Acta 2006, 52, 278. (36) Yakovlev, S.; Kharton, V.; Yaremchenko, A.; Kovalevsky, A.; Naumovich, E.; Frade, J. J. Eur. Ceram. Soc. 2007, 27, 4279. (37) Woodward, D. I.; Reaneya, I. M. Acta Crystallogr., Sect. B: Struct. Sci. 2005, 61. (38) Mather, G. C.; Garcia-Martin, S.; Benne, D.; Ritter, C.; Amador, U. J. Mater. Chem. 2011, 21, 5764. (39) Rodríguez, E.; Á lvarez, I.; López, M. L.; Veiga, M. L.; Pico, C. J. Solid State Chem. 1999, 148, 479. (40) Glazer, A. M. Acta Crystallogr., Sect. B: Struct. Sci. 1972, 28, 3384. (41) Howard, C. J.; Zhang, Z. J. Phys.: Condens. Matter 2003, 15, 4543. (42) Howard, C. J.; Zhang, Z. Acta Crystallogr., Sect. B: Struct. Sci. 2004, 60, 249. (43) Zhang, Z.; Howard, C. J.; Knight, K. S.; Lumpkin, G. R. Acta Crystallogr., Sect. B: Struct. Sci. 2006, 62, 60. (44) Garcia-Martin, S.; Alario-Franco, M. A.; Ehrenberg, H.; Rodriguez-Carvajal, J.; Amador, U. J. Am. Chem. Soc. 2004, 126, 3587. (45) Mitchell, R. H. Perovskite: Modern and Ancient; Almay Press, Inc.: Ontario, Canada, 2002. (46) Knapp, M. C.; Woodward, P. M. J. Solid State Chem. 2006, 179, 1076. (47) King, G.; Thimmaiah, S.; Dwivedi, A.; Woodward, P. M. Chem. Mater. 2007, 19, 6451. (48) Rodriguez, E.; Lopez, M. L.; Campo, J.; Veiga, M. L.; Pico, C. J. Mater. Chem. 2002, 12, 2798. (49) Solid State Physics; Ashcroft, N. W., Mermin, N. D., Eds.; Harcourt College Publishers: Fort Worth, TX, 1976. (50) Rampling, M. J.; Mather, G. C.; Marques, F. M. B.; Sinclair, D. C. J. Eur. Ceram. Soc. 2003, 23, 1911. (51) Pérez-Flores, J. C.; Ritter, C.; Pérez-Coll, D.; Mather, G. C.; Canales-Vázquez, J.; Gálvez-Sánchez, M.; García-Alvarado, F.; Amador, U. Int. J. Hydrogen Energy 2012, 37, 7242.

respectively, heated at 1373 K for 12 h in air. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*Fax: (34) 91 351 04 96. E-mail: [email protected] (J.C.P.F.), [email protected] (U.A). Author Contributions

The authors contributed equally. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank the Ministerio de Economiá y Competitividad (MINECO) and Comunidad de Madrid for funding the projects (Nos. MAT2010-19837-C06, PIB2010JP-00181, and S2009/PPQ-1626, respectively). Financial support from Universidad San Pablo is also acknowledged. D.P.C. is also grateful to the MINECO and CSIC for a Ramón y Cajal contract. Access to the neutron facilities at the Institut Laue Langevin is gratefully acknowledged.

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ABBREVIATIONS SOFC, solid oxide fuel cell; IT-SOFC, intermediate-temperature solid oxide fuel cell; XRD, X-ray diffraction; NPD, neutron powder diffraction; SAED, selected-area electron diffraction; TEM, transmission electron microscopy; HRTEM, high-resolution transmission electron microscopy; EDS, energy-dispersive X-ray spectroscopy; TGA, thermogravimetric analysis; SEM, scanning electron microscopy; ITF, isotropic thermal factor; FFT, fast Fourier transform; BVS, bond valence sum; AF, antiferromagnetic; SO, spin only; YSZ, yttriastabilized zirconia; LSM, La0.8Sr0.2MnO3; MIEC, mixed ionic−electronic conductivity

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REFERENCES

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A- and B-Site Ordering in the A-Cation-Deficient Perovskite Series La2

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