A-site deficiency, phase purity and crystal structure in lanthanum strontium ferrite powders

Solid State Ionics 178 (2007) 1326 – 1336 www.elsevier.com/locate/ssi

A-site deficiency, phase purity and crystal structure in lanthanum strontium ferrite powders T. Striker a,⁎, J.A. Ruud a , Y. Gao a , W.J. Heward a , C. Steinbruchel b a

b

GE Global Research Center, One Research Circle, Niskayuna, NY 12309, United States Department of Materials Science and Engineering, Rensselaer Polytechnic Institute, Troy, NY 12180, United States Received 25 March 2007; received in revised form 18 June 2007; accepted 18 June 2007

Abstract Lanthanum strontium ferrite (LSF) powders of average composition of (La1−x′ Srx′) y′FeO3+δ, were fabricated over a range of average strontium composition between 0.2 ≤ x′ ≤ 0.5 and average A-site occupancy between 0.8 ≤ y′ ≤ 1.0. Samples that were originally determined to be perovskite phase pure using X-ray diffraction with CuKα radiation were found to have significant amounts of non-perovskite phases when evaluated using high-energy synchrotron radiation. As-fabricated samples with nominal A-site deficiency, y′ b 1, contained a hexaferrite phase. When treated at 955 °C in pO2 = 10− 10 atm, those samples contained magnetite. The actual composition of the perovskite phase was corrected to account for the presence of the second phases through mass balance calculations. As a result, the actual A-site deficiency of the perovskite phase was much lower than the average value of the bulk powder. For as-fabricated powders with x′ b 0.4, it was determined that the A-site deficient LSF perovskite phases were metastable. At equilibrium, a mixture of A-site stoichiometric perovskite and hexaferrite phases was present. The refined perovskite crystal structures and unit cell volumes were consistent with literature trends. © 2007 Elsevier B.V. All rights reserved. Keywords: LSF; SOFC cathode; Perovskite; A-site stoichiometry; Phase analysis

1. Introduction Solid oxide fuel cells (SOFCs) offer the potential for highly efficient power generation systems. Advances in cell materials and microstructures for high specific power density and reduced operating temperature have been pursued on the path to the realization of the technology [1]. Of particular interest, is the development of new cathode materials as an alternative to the standard lanthanum strontium manganate (LSM)/yttria stabilized zirconia (YSZ) composite cathode [2,3]. For efficient oxygen reduction, SOFC cathode materials must be catalytically active and good conductors of electrons and ions. Lanthanum strontium ferrite (LSF) [4,5] and lanthanum strontium cobalt ferrite (LSCF) [6–8] are mixed ionic-electronic conducting (MIEC) perovskites in the ferrite family, and have shown promise

⁎ Corresponding author. GE Global Research, One Research Cicle, K-1 MB261, Niskayuna, NY 12309, United States. Tel.: +1 518 387 4352; fax: +1 518 387 6204. E-mail address: [email protected] (T. Striker). 0167-2738/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2007.06.011

as cathode materials by exhibiting higher catalytic activity and ionic conductivity than the LSM/YSZ composite. The perovskite crystal structure, ABO3, is well-known for the wide range of A-site and B-site elements which form stable compounds [9,10]. Substitution of divalent ions for trivalent ions on the A-site can be accommodated in the pervoskite structure by a change in oxygen stoichiometry or valence of the B-site ion. The LSF perovskite phase, (La1−xSrx)yFeO3+δ, is stable over a range of Sr contents, x, and oxygen sub-stoichiometry values, indicated with a negative δ value [11]. In a similar manner, the perovskite structure can also accommodate some amount of sub-stoichiometry on the A-site, y b 1. A-site deficiency is of particular interest for SOFC cathodes because of the potential for increasing the amount of oxygen vacancies, which results in an increased oxygen ion conductivity and perhaps also an increase in catalytic activity [12,13]. A-site non-stoichiometry has been explored for perovskites of interest to SOFCs, including manganites, ferrites and cobaltite–ferrites. For LSM, A-site non-stoichiometry has been reported for a wide range of compositions, 0.10 ≤x ≤ 0.70 and 0.90 ≤y ≤ 1.10,

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where (La1−xSrx)yMnO3+δ [14]. However, a modest amount of non-perovskite phases were observed even for stoichiometric Asite samples (as much as 13 wt.% of La2O3 at x = 0.35). Singlephase A-site deficient (La0.89Sr0.11)yMnO3+δ was synthesized for 0.89 b y b 0.94 [15]. In that study, a Mn2O3 secondary phase was observed for lower A-site occupancy levels, while a La2O3 phase was observed for higher values of y. A-site deficient LSF has also been reported. Oxygen diffusion coefficients were measured for single-phase LSF, with x = 0.25 and y = 0.90 [16]. In addition, electrochemical performance was evaluated for (La0.8Sr0.2)yFeO3+δ cathodes with 0.95≤y ≤ 1.05 [4]. The A-site sub-stoichiometric compositions had faster sintering kinetics than the stoichoimetric composition. For cathodes sintered at the same temperature, power density was highest for the stoichiometric cathodes, but the microstructures were not optimized for the differences in sintering kinetics. A-site deficiencies as high as 20% were observed in (La0.8Sr0.2)yFeO3+δ powders and reported to be single-phase [17]. A comparison of electrochemical performance of A-site deficient LSF and LSCF cathodes has been made [18]. Phase pure powders with 1% A-site deficiency in LSF and LSCF were produced for x = 0.2 and 0.4. Sintering kinetics were enhanced with A-site deficient LSCF but not for LSF. For (La0.6Sr0.4)Co0.2Fe0.8O3+δ, A-site deficiency of 1% increased the electrochemical performance of the SOFCs; however, deficiencies greater than 5% had a negative effect on measured performance. Single-phase LSCF materials of a similar composition with A-site deficiencies up to 27% have also been reported [19,6]. Although A-site deficient ferrite based perovskites have been evaluated and used for SOFC cathodes, there have been few systematic studies of the influence of A-site deficiency levels on phase purity and phase structure. The crystal structure [11] and oxygen stoichiometry [20] of A-site stoichiometric LSF have been evaluated for a range of Sr contents and oxygen partial pressures. The objective of this work was to produce LSF materials with an average composition of (La1−x′Srx′)y′FeO3+δ, fabricated over a strontium composition range of 0.2≤x′ ≤ 0.5 and an A-site occupancy range of 0.8 ≤y′ ≤ 1.0. Phase purity and crystal structure of the powders were evaluated using synchrotron X-ray diffraction, and oxygen stoichiometry was measured using a redox technique. The observed phase purity and the implications for the stability of A-site deficient LSF phases are discussed. 2. Experimental Powders with an average composition of (La1−x′Srx′)y′FeO3+δ were fabricated using a glycine–nitrate synthesis [21]. The average strontium composition on the A-site, x′, is equal to XSr / (XSr +XLa), where XSr and XLa are the molar concentrations of strontium and lanthanum respectively. The average A-site occupation, y′, is equal to (XLa +XSr) /XFe, where XFe is the molar concentration of iron. A Varian Liberty II inductively coupled plasma atomic emission spectrometer (ICP-AES) was used to analyze the molarities of stock solutions prepared using lanthanum, strontium and iron nitrates. To produce the desired compositions, appropriate amounts of nitrate stock solutions were mixed and combined with glycine (using 5 mol% glycine excess). The solution was

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dried and heated in a stainless steel beaker with a mesh chimney to approximately 200 °C, where the glycine–nitrate solution combusted producing a fine ash. The resulting ash was calcined at 850 °C for six hours, followed by subsequent heat treatments at 1000 °C and 1200 °C for six hours with intermittent grindings between each heat treatment. Average strontium contents, x′, ranged from 0 to 0.5, with most concentrations between 0.2 and 0.4. Average A-site occupations, y′, ranged from 0.77 to 1.03, with most concentrations between 0.9 and 1.0. Twenty-five powders were prepared in-house and five powders were obtained from Praxair Specialty Ceramics (Woodinville, WA). The actual compositions of the as-fabricated powders were measured using ICP-AES, typically with a standard deviation of less than 0.005 and 0.011 for x′ and y′, respectively. LSF powders were heat-treated to modify the oxygen stoichiometry. The powders were heated inside an alumina tube, with a thermocouple attached to monitor temperature. The first heat treatment was at 955 °C for 15 h under a flowing mixture of CO and CO2 of a ratio to produce a pO2 of 10− 10 atm. The heat treatment objective was to reduce the perovskite powders so that the average valence state of the Fe was 3+ [20]. Following the 15 hour hold, the powders were quenched from the hot-zone while maintaining desired oxygen partial pressures, typically cooling the samples to less than 450 °C in 10 min. A second heat treatment in pO2 = 10− 5 atm was done at 800 °C under a flow of 10 ppm oxygen (balance argon), held for 15 h and quenched in a similar manner. The perovskite oxygen stoichiometry was determined by measuring the amounts of trivalent iron present using a redox technique [22]. Approximately 0.1 g of powder was weighed and mixed with 10 mL of a 0.1 M iron sulfate solution. Around 200 mL of sulfuric acid (50% H2SO4 and 50% H2O) was then added and heated while stirring. The powder dissolved in the acid to produce a solution of cations, which included Fe3+ and Fe4+. Excess ferrous sulfate solution was added to convert the Fe4+ to Fe3+, then a 0.1 N potassium permanganate was titrated into the solution. Total Fe2+ consumption was signified as the clear solution turned pink, indicating the amount of Fe4+ that was in the perovskite powder. The average valence state of iron in the (La1−xSrx)yFe4+1−aFe3+aO3+δ perovskite powders was calculated, where a is the amount of Fe3+ present. Then, using charge balance, the oxygen stoichiometry was determined. In order to verify that the quenching procedure was satisfactory and that the titration measurement was accurate, a comparison with the literature was performed. Two A-site stoichiometric samples of composition x′ = 0.2 and 0.45 were heat-treated at two extreme levels of pO2 (10− 10 and 1 atm) and quenched. The oxygen stoichiometries determined were in good agreement with the results from Mizusaki et al. [20]. This demonstrated that the samples were quenched at a sufficient rate to avoid oxygen absorption upon cooling, and that the oxygen stoichiometry was measured with an acceptable error. The powders were initially characterized with X-ray powder diffraction (XRD) in-house with a Bruker DB Advance diffractometer using CuKα radiation, and an MBraun PSD-50 linear position-sensitive-detector. The step size was 0.0066°, and a 0.2 s/step count time was used. To optimize detection of nonperovskite phases by reducing sample fluorescence, a graphite

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Fig. 1. CuKα XRD patterns for sample LS43F-95 calcined at 1000 °C and 1200 °C. Bragg reflections are shown for (La0.4Sr0.6)FeO3 (top), SrLaFeO4 (middle) and SrFe12O19 (bottom). SrLaFeO4 is present in the sample calcined at 1000 °C.

diffracted beam monochromator coupled with a scintillating detector and an energy dispersive detector were used. An extensive set of X-ray powder diffraction measurements was made using high-energy (60–80 keV) synchrotron X-rays from the X17B1 beamline at the National Synchrotron Light Source at Brookhaven National Laboratory. A 4.2 T X17 wiggler provided a high flux of collimated high-energy X-rays, with a critical energy of 22 keV. A monochromatic X-ray beam of energyresolution of 10− 4 (dE /E) was provided by using a sagittalfocusing silicon double-crystal monochromator, with both crystals in asymmetric laue mode [23]. The beam size, defined by two pairs of horizontal (0.5 mm) and vertical (0.1 mm) slits, was focused on the sample with a horizontal divergence of 50 μrad, and vertical divergence between 2 and 4 μrad. The diffraction plane was

oriented vertically to take advantage of the low vertical divergence of the beam, resulting in better angular resolution. The measurements were carried out in transmission mode, using an automated, multiple-sample holder to achieve high measurement throughput. Powder samples were loaded into holes (∼3 mm in diameter) on a ∼1 mm thick aluminum plate, and were sealed with Kapton film on both sides. NIST CeO2 powder was used as an external standard for achieving lattice parameter accuracy. Diffraction data was collected using MarCCD (MarUSA, Inc.), a two-dimensional detector capable of rapid data acquisition. Typical data collection time was 2–3 min per sample. The 2D data was integrated in the diffraction plane using a custom-developed program written on the IDL platform (Research Systems, Inc.) [24]. The combined use of high-energy synchrotron X-rays and a 2D detector provided a

Fig. 2. Synchrotron XRD pattern for LS43F-95 calcined at 1200 °C. Experimental (+) and calculated (−) data are plotted on the top with a difference trace (experimental minus calculated data) plotted on the bottom. Vertical bars in the middle represent SrFe12O19 (top tick marks), rhombohedral perovskite (middle tick marks) and orthorhombic perovskite (bottom tick marks) Bragg reflections.

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composition of the perovskite phase, (La1−xSrx)yFeO3+δ, was determined from the average bulk composition by mass balance when considering the non-perovskite phases present. For singlephase powders, the composition of the perovskite phase is the same as the bulk powder composition, i.e., x =x′ and y =y′. 3. Results

Fig. 3. a) CuKα XRD pattern compared with b) synchrotron XRD pattern for sample LS53F-84. Positions of Bragg reflections are indicated for the perovskite phase (+) and SrFe12O19 (●,○) are labeled. 6 volume percent of SrFe12O19 is apparent in the XRD pattern taken with synchrotron radiation but not with CuKα radiation.

practical alternative to in-house instrumentation for achieving high-angular resolution, high-sensitivity to secondary phases, and high-throughput for mapping a large compositional space. Rietveld refinement [25] of the patterns was completed using the General Structure Analysis System (GSAS) software [26]. The refinement parameters generally included lattice parameters, atomic positions, scale factors, and variables for background and peak-shape functions. The initial crystal structure parameters for ¯c), and cubic (pm3m) the orthorhombic (pbnm), rhombohedral (r3 perovskite phases were obtained from literature [11]. Significant amounts of secondary phases were detected in the synchrotron X-ray measurements. The refinement parameters (lattice sizes and atomic positions) for SrFe12O19 [27], SrLaFeO4 [28], Fe3O4 [29], and Sr2LaFe3O8 [30] were fixed at values reported in literature. Relative phase compositions were determined from XRD pattern refinement. The figure-of-merit parameters, Rp, Rwp, and χ2 (typically less than 0.1, 0.15 for Rp and Rwp, respectively, and less than 5 for χ2) were monitored to ensure satisfactory agreement between observed and calculated diffraction data. The actual

A calcination procedure was developed to produce LSF powders that were nominally single-phase as determined by CuKα radiation powder XRD. Two powders with large strontium contents (x′ ∼0.3 and 0.4) and nominal A-site deficiency were synthesized and calcined sequentially for 6 h at 850 °C, 1000 °C and 1200 °C. At the lower calcination temperatures, SrLaFeO4 phase was present, as shown in Fig. 1 for sample LS43F-95. However, after the 1200 °C heat treatment, this phase was no longer detected. Based on those results, all powders in the study received the same calcination treatment up to 6 h at 1200 °C. The five vendor powders were evaluated in the as-received state and verified to be nominally phase pure using the CuKα XRD measurements. It was found that substantial amounts of a second, nonperovskite, phase were observed in the synchrotron diffraction data that were not detectable using CuKα XRD. The synchrotron pattern for sample LS43F-95 calcined at 1200 °C is shown in Fig. 2. The synchrotron pattern clearly shows the presence of SrFe12O19 with strong reflections at approximately 3.22°, 3.62° and 3.93° (2.94 Å, 2.62 Å, and 2.42 Å, respectively) that have the expected relative intensities. A more detailed comparison between CuKα and synchrotron data for sample LS52F-84 is shown in Fig. 3. Though lying on different 2θ scales, the patterns have been overlaid in the same regions to show the main peaks of interest. No clear evidence of the SrFe12O19 phase was observed using the CuKα XRD. The main SrFe12O19 peaks, (107) with 87% intensity and (114) with 100% intensity, are confounded by the perovskite orthorhombic (112) and (021) reflections, respectively (around 32.3° and 34.2° two-theta on the CuKα pattern), and can therefore be difficult to resolve. Secondary SrFe12O19 peaks with about the same intensity should be at 30.4° and 37.2°. Since they are not evident, it is difficult to conclude that the SrFe12O19 is present using standard CuKα XRD techniques. The same powder shows strong evidence of the SrFe12O19 phase when measured using high-energy synchrotron X-rays. This indicates that an intense Xray source is needed for reliable phase analysis of nominally A-site deficient LSF powders. The phase contents of the as-fabricated powders (including vendor powders) using synchrotron XRD are reported in Table 1. Powders with average bulk A-site occupation, y′ N 1 had the secondary phase SrLaFeO4, and those with y′ b 1 had the secondary phase SrFe12O19. As y′ decreased, the concentration of SrFe12O19 increased, as shown in Fig. 4. As much as 15 vol.% SrFe12O19 was observed for the powder with the lowest y′ value. As-fabricated powders with higher Sr contents generally had lower amounts of SrFe12O19. For instance, the powder with y′ = 0.82 and x′ = 0.4 showed approximately half of the SrFe12O19 phase as the samples with the same A-site occupation but with Sr contents of x′ = 0.3 and 0.2.

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Table 1 Volume fraction of phases and corrected perovskite compositions for as-fabricated powders Bulk composition

Volume fraction phase

Corrected perovskite composition

Oxygen deficiency

Sample ID

x′

y′

Perovskite

SrFe12O19

SrLaFeO4

x

y

δ

LS20F-01-V LS20F-98 LS22F-94 LS20F-93-V LS22F-89-V LS19F-82 LS29F-03 LS30F-02 LS32F-98 LS31F-95-V LS27F-88-V LS33F-82 LS28F-77 LS40F-01 LS39F-98 LS43F-95 LS40F-91 LS39F-80 LS45F-00 LS53F-96 LS52F-84

0.20 0.20 0.22 0.20 0.22 0.19 0.30 0.30 0.32 0.31 0.27 0.33 0.28 0.40 0.39 0.43 0.40 0.39 0.45 0.53 0.52

1.01 0.98 0.94 0.93 0.89 0.82 1.03 1.02 0.98 0.95 0.88 0.82 0.77 1.01 0.98 0.95 0.91 0.80 1.00 0.96 0.84

1.000 0.983 0.958 0.974 0.938 0.892 0.924 1.000 0.978 0.959 0.958 0.872 0.852 0.964 1.000 0.978 0.968 0.945 1.000 1.000 0.940

0.000 0.017 0.042 0.026 0.062 0.108 0.000 0.000 0.022 0.041 0.042 0.128 0.148 0.000 0.000 0.022 0.032 0.055 0.000 0.000 0.060

0.000 0.000 0.000 0.000 0.000 0.000 0.076 0.000 0.000 0.000 0.000 0.000 0.000 0.036 0.000 0.000 0.000 0.000 0.000 0.000 0.000

0.20 0.20 0.21 0.19 0.21 0.17 0.28 0.30 0.31 0.31 0.26 +/− 0.005 0.31 0.26 0.40 0.39 0.43 0.40 +/− 0.007 0.38 +/− 0.007 0.45 0.53 +/− 0.003 0.52 +/− 0.007

1.01 1.02 1.02 0.98 1.01 1.03 0.98 1.02 1.02 1.00 0.95 +/− 0.012 1.00 1.03 0.99 0.98 0.99 0.95 +/− 0.012 0.88 +/− 0.011 1.00 0.96 +/− 0.006 0.94 +/− 0.010

0.01 0.03 0.01 − 0.08 − 0.04 − 0.03 − 0.05 0.02 − 0.07 − 0.02 − 0.10 − 0.00 0.03 − 0.08 − 0.05 − 0.05 − 0.10 − 0.19 − 0.02 − 0.12 − 0.11

Samples with the designation -V were obtained from a commercial vendor. Standard deviations are listed for the five statistically significant A-site deficient perovskites.

SrLaFeO4 was observed in two samples with y′ N 1, showing increasing concentration with increasing y′. The phase contents of the powders heat-treated at 800 °C, with pO2 = 10− 5 atm are reported in Table 2. The amounts of SrFe12O19 and LaSrFeO4 were consistent with as-fabricated powders. Phase contents of powders heat-treated at 955 °C and pO2 = 10− 10 atm are reported in Table 3. Secondary phases, Fe3O4 and Sr2LaFe3O8, were observed in addition to the perovskite phases. As y′ decreased, increasing amounts of Fe3O4 were observed. Approximately 12 vol.% Fe3O4 was observed at the highest bulk A-site deficiency. Another non-perovskite phase, Sr2LaFe3O8, was observed for powders with x′ ≥ 0.45. For powders with secondary phases present, the perovskite phase composition differed from the bulk powder composition. The x and y values in the perovskite were higher than the corresponding x′ and y′ bulk sample concentrations, as shown in Table 1. Only five of the as-fabricated perovskites had statistically significant levels of A-site deficiency (samples LS27F-88-V, LS40F-91, LS39F-80, LS53F-96 and LS52F-84) as indicated by the standard deviations for these samples. Typical standard deviations for x and y were around 0.005 and 0.012, respectively. The largest observed perovskite A-site deficiency was y = 0.88 (for x = 0.38). Titrations of the as-fabricated and heat-treated powders were used to determine the average iron valence state in the perovskite phases, using the perovskite x and y values corrected for nonperovskite phases present. Using charge balance, the oxygen stoichiometry was determined for the as-fabricated and the heattreated perovskites at two oxygen partial pressures. Perovskite oxygen stoichiometries, δ, were determined for the as-fabricated powders. In general, as the perovskite A-site

deficiency increased, the oxygen stoichiometry decreased, as shown in Fig. 5. A minimum δ value of approximately −0.19 was observed for the sample with y = 0.88. Sr content, x, did not seem to have a significant effect on δ values for the as-fabricated powders. The oxygen stoichiometries for the perovskite phase in powders heat-treated at 955 °C at pO2 = 10− 10 atm are shown in Fig. 6. Oxygen stoichiometry decreased as A-site deficiency increased for powders of varying strontium content. The predicted

Fig. 4. Volume fraction of SrFe12O19 in as-fabricated powders as a function of average A-site occupation, y′ for powders with average Sr contents of x′ ∼0.2 (+), 0.3 (○) and 0.4 ( ). The top and bottom lines are added for x′ ∼0.2 and 0.4 respectively as a visual guide.

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Table 2 Volume fraction of phases and corrected perovskite compositions for powders heat-treated at 800 °C in pO2 = 10− 5 atm Bulk composition

Volume fraction phase

Corrected perovskite composition

Oxygen sub-stoichiometry

Sample ID

x′

y′

Perovskite

SrFe12O19

SrLaFeO4

x

y

δ

LS20F-01-V LS20F-98 LS22F-94 LS20F-93-V LS32F-98 LS31F-95-V LS40F-01 LS39F-98 LS43F-95

0.20 0.20 0.22 0.20 0.32 0.31 0.40 0.39 0.43

1.01 0.98 0.94 0.93 0.98 0.95 1.01 0.98 0.95

1.000 0.970 0.960 0.972 0.957 0.988 0.954 0.988 0.971

0.000 0.030 0.040 0.028 0.043 0.012 0.000 0.012 0.029

0.000 0.000 0.000 0.000 0.000 0.000 0.046 0.000 0.000

0.20 0.20 0.22 0.19 0.32 0.31 0.40 0.39 0.43

1.01 1.00 1.00 0.98 1.00 0.98 0.99 1.00 1.00

0.00 0.00 − 0.02 − 0.01 − 0.05 − 0.03 − 0.10 − 0.09 − 0.12

Samples with the designation -V were obtained from a commercial vendor.

oxygen deficiencies of the reduced powders, assuming the average Fe valence was +3 [20], are also plotted for reference. Some error exists between the experimental data and the predicted data. This is most likely due to the propagated error of the mass balanced perovskite compositions (y). The crystal structures for the perovskites in the as-fabricated powders were observed to be orthorhombic (pbnm) or a mixture of ¯c) crystal structures as shown by the pbnm and rhombohedral (r3 crystal structure ratios plotted in Fig. 7. These results are reported under the assumption that the perovskite stoichiometry differences are negligible when multiple perovskite crystal structures are observed. For strontium contents below 0.2, the orthorhombic phase (pbnm) was more prevalent. A transition at about x = 0.2 occurred where the amount of orthorhombic decreased, which increased the volume fraction of the rhombohedral perovskite phase. A similar trend was observed for heat-treated powders at 800 °C in pO2 = 10− 5 atm. Heat-treated powders at 955 °C under

pO2 = 10− 10 atm showed a mixture of orthorhombic and cubic perovskite phases as shown in Fig. 8. An orthorhombic-to-cubic transition occurred at x = 0.2, analogous to the orthorhombic-torhombohedral transition seen at higher oxygen partial pressures. Vendor powder perovskite crystal structures were found to be consistent with the perovskite structures fabricated in-house. In addition, the perovskite phase transitions are similar to those observed in prior work [31,32]. Lattice parameters were also determined for each perovskite phase, including unit cell volume and gram-atomic cell volume in cubic angstroms. The unit cell volumes were converted to gramatomic cell volumes by dividing by the formula units per unit cell, Z. For cubic, orthorhombic, and rhombohedral perovskite crystal structures, Z is 1, 4, and 6, respectively. In general, an increase in the strontium content of the as-fabricated perovskites correlated to a decrease in the gram-atomic cell volumes as shown in Fig. 9, which is consistent with the literature [11]. A comparison of the

Table 3 The volume fraction of phases and corrected perovskite compositions for 955 °C heat-treated powders in pO2 = 10− 10 atm Bulk composition

Volume fraction phase

Corrected perovskite composition

Oxygen sub-stoichiometry

Sample ID

x′

y′

Perovskite

Fe3O4

Sr2LaFe3O8

x

y

δ

LS20F-01-V LS20F-98 LS22F-94 LS19F-82 LS30F-03 LS32F-98 LS31F-95-V LS27F-88-V LS33F-82 LS40F-01 LS39F-98 LS43F-95 LS40F-91 LS39F-80 LS45F-00 LS52F-84 LS53F-96

0.20 0.20 0.22 0.19 0.30 0.32 0.31 0.27 0.33 0.40 0.39 0.43 0.40 0.39 0.45 0.52 0.53

1.01 0.98 0.94 0.82 1.03 0.98 0.95 0.88 0.82 1.01 0.98 0.95 0.91 0.80 1.00 0.84 0.96

1.000 0.991 0.980 0.886 1.000 0.982 0.999 0.961 0.934 1.000 0.994 0.987 0.958 0.905 0.781 0.673 0.726

0.000 0.009 0.020 0.114 0.000 0.018 0.001 0.039 0.066 0.000 0.006 0.013 0.042 0.095 0.000 0.050 0.000

0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.000 0.219 0.277 0.274

0.20 0.20 0.22 0.19 0.30 0.32 0.32 0.27 0.33 0.40 0.39 0.43 0.40 0.39 0.37 0.43 0.45

1.01 1.00 0.99 1.00 1.03 1.02 0.96 0.97 0.97 1.00 0.98 0.99 1.00 1.00 1.00 0.92 0.95

− 0.08 − 0.08 − 0.10 − 0.07 − 0.13 − 0.10 − 0.17 − 0.17 − 0.21 − 0.15 − 0.23 − 0.20 − 0.20 − 0.21 − 0.19 − 0.33 − 0.30

Samples with the designation -V were obtained from a commercial vendor.

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Fig. 5. Oxygen deficiency (δ) versus A-site occupation (y) for as-fabricated perovskite phases, where (La1−xSrx)yFeO3+δ. Sr contents of x′ ∼ 0.2 (×), 0.3 (○) and 0.4 ( ) are noted. Error bars represent one standard deviation.

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weighted average gram-atomic cell volumes (weighted for volume fraction of each phase) for each oxygen partial pressure heat treatment is also included. Lower pO2 treatments produced higher weighted average gram-atomic cell volumes. 4. Discussion It is generally thought that perovskites are tolerant to some A-site sub-stoichiometry without the presence of significant amounts of non-perovskite secondary phases. A-site deficient perovskite phases have been reported for LSM [14–16,33], LSF

Fig. 6. Oxygen deficiency (δ) versus A-site occupation (y) for perovskite phases in powders heat-treated at 955 °C, pO2 = 10− 10 atm. Sr contents of x′ ∼ 0.2 (×), 0.3 (○) and 0.4 ( ) are noted. Lines indicate the calculated δ for x = 0.2, 0.3 and 0.4 (assuming Fe valence is 3+). The error bars correspond to one standard deviation.

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Fig. 7. Volume fraction of orthorhombic and rhombohedral phases in the asfabricated perovskite structure as a function of Sr content. A transition around x =0.2 is apparent, where the orthorhombic perovskite phase decreases in favor of more rhombohedral phase.

[4,8,16,17], and LSCF [6,8,19]. In most cases the perovskite phases were determined to be phase pure by CuKα XRD or small amounts of second phases were present which were not considered detrimental to performance. However, in the present study it was observed that for A-site deficient lanthanum strontium ferrite, significant amounts of a hexaferrite phase were present in both in-house and vendor produced powders. The hexaferrite phase was not easily resolvable by standard CuKα XRD, yet it had to be quantitatively included in the mass balance in order for an accurate determination of the actual stoichiometry of the perovskite phase. Of the 13 LSF powders fabricated with apparent A-site deficiency, only 5 had perovskite phases with significant A-site deficiency after considering the mass balance of the non-perovskite phases.

Fig. 8. Volume fraction of orthorhombic and cubic perovskite phases powders heat-treated at 955 °C in pO2 = 10− 10 atm. A transition around x = 0.2 is apparent, where the orthorhombic perovskite phase decreases in favor of more cubic.

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Fig. 9. Weighted gram-atomic cell volumes as a function of oxygen partial pressure and Sr content, x., for as-fabricated powders (×) and powders heat-treated at pO2 =10− 5 atm (○) and pO2 = 10− 10 atm ( ). Linear regression lines are shown.

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The high-energy synchrotron radiation allowed for larger powder volumes to be examined and produced higher signalto-noise measurements, enabling the detection of the hexaferrite phase. Because the primary hexaferrite XRD peaks overlapped with the main perovskite peaks, the hexaferrite phase, in particular, required synchrotron XRD for its detection in LSF powders. The phase diagram for La2O 3–SrO–Fe 2O3 indicates the existence of a strontium lanthanum hexaferrite phase (Sr1−zLazFe12O19) [34] at temperatures of 1100 °C and above. The solubility of La in the strontium hexaferrite phase is low at 1100 °C and it increases with temperature. Pure LaFe12O19 stability is limited to temperatures between 1380 °C and 1420 °C [35]. Since only small amounts of La solubility are predicted in the compositional ranges of interest for this study (less than 30% La on the Sr site based on a linear

Fig. 10. A predicted isothermal section of the La2O3–SrO–Fe2O3 phase diagram at 1200 °C, based on prior work [34]. Samples LS19F-82 and LS33F-82 are shown for average bulk compositions (closed symbols) and the average perovskite compositions (open symbols) with the presence of hexaferrite phase. The average bulk composition of LS39F-80 (●) is compared to the as-fabricated (○) and the 60 h heat-treated (⊗) perovskite compositions.

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interpolation), the hexaferrite phase was assumed to be stoichiometric SrFe12O19. Since hexaferrite is iron rich, ignoring La contribution to mass balancing corrections of the perovskite compositions should have a very minimal effect (an approximate 5% difference in x). As a result of the presence of the hexaferrite phase in the Asite deficient bulk powders, the average A-site occupancy of the perovskite phase was larger than the A-site occupation expected from the average bulk composition (y N y′). The nominal sample compositions can be mapped on a schematic of the Fe2O3-rich section of the ternary phase diagram at 1200 °C (which is consistent with published phase diagrams at 1100 °C and 1300 °C [34]) as shown in Fig. 10. Samples with y′ of approximately 0.80 and three levels of x′ are plotted. All three samples fall in the two-phase La1−xSrxFeO3+δ–Sr1−zLazFe12O19 region. Two as-fabricated samples (LS19F-82 and LS33F-82) with lower Sr content showed the presence of a hexaferrite phase and A-site stoichiometric perovskite compositions, as plotted on the LSF solid solution line. The as-fabricated sample (LS39F-80) with higher Sr content, x′ = 0.39, also showed a hexaferrite phase (1.1 mol%). However, when adjusting for mass balance, the average A-site occupation of the perovskite was 0.88. Upon further heat treatment of that sample (60 h at 1200 °C), 2.7 mol% hexaferrite was observed and the corrected perovskite composition fell on the solid solution line, representing an A-site stoichiometric LSF phase. The same 60 hour heat treatment at 1200 °C for samples of composition x′ b 0.4 showed that as-fabricated A-site deficiency was metastable. After the 60 hour heat treatment, all x′ b 0.4 samples exhibited an A-site stoichiometric LSF perovskite phase, demonstrating that A-site deficiency is not stable for lower Sr containing samples. However, the largest A-site deficiency in an as-fabricated phase pure perovskite powder was observed for sample LS53F-96 (where x′ N 0.4). This powder was not subjected to the additional heat treatment; therefore, the possibility of stable 0.95 b y b 1.0 cannot be ruled out for high Sr contents.

Fig. 11. SEM micrograph of LS39F-80 heat-treated for 60 h at 1200 °C. The perovskite and hexaferrite phases are the light and dark gray regions, respectively. The black regions are porosity.

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Fig. 12. Proposed mechanisms for hexaferrite formation illustrated with LS39F-80 sample. In the first mechanism (a, top), perovskite composition, x and y, changes as hexaferrite grows. In the second (b, bottom), perovskite x and y (x′ and y′ for the non-stoichiometric perovskite) remain constant, and the amount of A-site stoichiometric perovskite grows in conjunction with the hexaferrite.

The microstructure of a pressed pellet of LS39F-80, after the 60 hour heat treatment at 1200 °C, is shown in Fig. 11. The backscatter scanning electron (BSE) micrograph shows two distinct phases. From energy dispersive X-ray spectroscopy (EDX), it was determined that the lighter regions represent the perovskite phase, the darker regions represent the iron rich hexaferrite phase, and the black areas are porosity. In addition to the large scale porosity from incomplete sintering of the pressed pellet, fine scale porosity is observed at the interfacial regions between the hexaferrite and perovskite phases in the microstructure, which could be attributed to the molar volume differences. The A-site deficient LSF perovskite is most likely a metastable phase, then, that evolves with time at temperature to form a mixture of an A-site stoichiometric perovskite and a hexaferrite phase. The glycine–nitrate synthesis used to fabricate the powder provides intimate mixing of the cations. For bulk compositions of y′ b 1, a metastable A-site deficient perovskite phase is crystallized during fabrication and the hexaferrite is nucleated during heat treatment. Two mechanistic approaches are proposed for the subsequent phase evolution using the LS39F-80 sample as a quantitative example. In the first mechanism, shown in Fig. 12a, the hexaferrite phase precipitates and grows in a uniform A-site deficient perovskite matrix. Initially, y = 0.8, and when the equilibrium amount of hexaferrite has been attained, y = 1. Growth of the hexaferrite phase involves the diffusion of La and Sr away from the hexaferrite phase interface into the A-site deficient perovskite matrix, thus increasing the La and Sr content of the perovskite. The intermediate state that was observed after the first heat treatment

schedule of 6 h at 1200 °C, was 1.1 mol% of hexaferrite in a uniform perovskite matrix of average composition y = 0.88 and x = 0.38. Further heat treatment and Sr diffusion continued the hexaferrite phase growth, and at equilibrium resulted in a final perovskite composition of y = 1.0 and x = 0.37, according to the following final mass balanced reaction: ðLa0:61 Sr0:39 Þ0:80 FeO3þδ →0:021SrFe12 O19 þ 0:75ðLa0:63 Sr0:37 ÞFeO3þδ

ð1Þ

An alternative mechanism is shown in Fig. 12b. In this case, when hexaferrite is nucleated initially, an adjacent A-site stoichiometric phase is also nucleated. The A-site stoichiometric phase grows concurrently, according to Eq. (1), while the metastable A-site deficient matrix phase of x′ = 0.39 and y′ = 0.80 remains at constant average composition. The hexaferrite interface provides a source for Sr diffusion, and the A-site stoichiometric interface provides a sink for Sr. In accordance with Eq. (1), the predicted phase compositions in the intermediate state (6 h at 1200 °C) represented by the as-fabricated LS39F-80 powder is 58.4 mol% x ′ = 0.39, y′ = 0.80 and 40.5 mol% x = 0.37, y = 1.0 for the 1.1 mol% of SrFe12O19 observed using refinement. With subsequent heat treatments, both the hexaferrite and A-site stoichiometric phases continue to grow until equilibrium is reached, when the final composition is 97.3 mol% x = 0.37, y = 1.0 and 2.7 mol% SrFe12O19. Because of peak convolution in the patterns of the perovskite phases, it was not possible to make a quantitative distinction using XRD measurements between the intermediate stages of either mechanism.

T. Striker et al. / Solid State Ionics 178 (2007) 1326–1336

Fig. 13. Projection of the unit cells for hexagonal SrFe12O19 (left) and rhombohedral (LaSr)FeO3 (right) on the (010) plane produced by what Diamond © 3.1 software.

A comparison of the unit cell structures observed in the asfabricated LS39F-80 sample, the hexagonal (P63/MMC) hexaferrite and the rhombohedral (R3¯c) perovskite, is shown in Fig. 13. The atomic positions of the Fe atoms are more similarly aligned for the two structures than the La and Sr positions. Two rhombohedral perovskite unit cells stacked in the c direction have similar lattice parameters to the hexaferrite unit cell, and therefore similar volumes. In the roughly equivalent structures, the number of Fe atoms is 24 and 12 for the hexaferrite and rhombohedral perovskite, respectively. In contrast, the number of A-site atoms (Sr and La) is 2 and 12 for the hexaferrite and rhombohedral perovskite, respectively. Therefore, when forming the hexaferrite phase, the amount of Fe must increase by a factor of two for roughly the same volume. However, for the same change, the amount of A-site atoms must decrease by a factor of six. Therefore, precipitation and growth of the hexaferrite phase may be limited by Sr and La diffusion from the perovskite phase into the A-site deficient perovskite matrix. As previously shown, an increase in the Sr content of the perovskite structure led to a decrease in the unit cell volume. Lower unit cell volumes may result in a decreased diffusion coefficient of Sr and La atoms in the perovskite structure. In addition, the larger Shannon radii of Sr [36] may result in a lower diffusion coefficient compared to La. For both cases, the rate of hexaferrite formation would be slower with increased Sr content. That may explain the more sluggish kinetics observed for hexaferrite formation in compositions with higher Sr content. The powders in this study are similar to those that have been studied for LSF SOFC cathode materials. Calcining and

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sintering temperatures for LSF powders are reported to be between 1000 °C–1200 °C, and between 1100 °C–1250 °C, respectively [4]; another study reports 1100 °C and 1150 °C, respectively [8]. The as-fabricated powders in this study were treated under similar conditions. Moreover, the vendor materials are typical of as-received powders that are often processed and deposited onto a SOFC electrolyte. Therefore, most nominally A-site deficient LSF powders used for SOFCs are expected to have some amount of hexaferrite present. The presence of up to 4 vol.% hexaferrite in both in-house and vendor samples of LSF with y′ = 0.95 is noteworthy and relevant to SOFC cathodes. Cathodes of similar y′ values have been studied using various experimental setups [4,8,17]. Cathode properties, such as conductivity or catalytic activity, attributed to A-site deficiency may be a consequence of the presence of the second phase. It is also possible that the second phase may enhance the densification of the cathode during sintering. The heat-treated powders at 800 °C in pO2 = 10− 5 atm are also relevant to SOFCs. During operation, as a result of electrochemical bias, LSF cathodes are expected to have operating oxygen partial pressures no lower than 10− 5 atm [37]. In the present study, the hexaferrite phase was not observed in LSF at pO2 = 10− 10 atm. However, hexaferrite was present in powders heat-treated at 800 °C in pO2 = 10− 5 atm, and in similar concentrations to those observed in the asfabricated powders. As a result, hexaferrite presence is also expected for LSF cathodes during SOFC operation. 5. Conclusions The deficiency of lanthanum and strontium, with respect to iron, in lanthanum strontium ferrite powder, (La1−x′Srx′)y′FeO3+δ, has a significant effect on perovskite phase purity. Powders that were originally determined to be phase pure perovskites using CuKα XRD were found to have non-perovskite phases present when analyzed using synchrotron XRD. Therefore, for accurate phase analysis of nominally A-site deficient LSF perovskites (y′ b 1), CuKα radiation appears insufficient. In high and medium oxygen partial pressures (as fabricated and pO2 = 10− 5 atm, respectively), most LSF powders with y′ b 1 contained a mixture of a hexaferrite phase and a perovskite phase. Because of the presence of the hexaferrite phase, the A-site deficiency of the perovskite phase was much less than the average A-site deficiency of the bulk powder composition. Even though LSF powders with an average A-site occupancy as low as y′ =0.77 were fabricated, the lowest A-site occupancy observed in the perovskite phase was y = 0.88. However, after further heat treatment, the perovskite phase in that powder became A-site stoichiometric with a composition of hexaferrite as predicted by the phase diagram. For all powders in the study with x′ b 0.4, A-site deficiency was not stable and additional heat treatment resulted in A-site stoichiometric perovskites. For powders heat-treated in pO2 = 10− 10 atm, Fe3O4 was observed for powder with y′ b 1. Perovskite crystal structures and unit cell volumes for the as-fabricated powders were consistent with literature trends. Namely, the orthorhombic phase transitioned to a rhombohedral perovskite crystal structure as strontium content

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increased. In pO2 = 10− 10 atm, the orthorhombic phase changed to the cubic crystal structure as the strontium content increased. Acknowledgements The authors acknowledge the support of Brookhaven National Laboratory for the use of the X17B1 beamline at the National Synchrotron Light Source. The work was supported by the U.S. Department of Energy under contract No. DE-AC0298CH10886 and the Brookhaven National Laboratory LDRD 05-057. A. Krushevska and A. Young provided elemental analysis. In addition, K. Luthra, S. Renou, P. Meschter, and B. Ramamurthi offered helpful technical guidance. References [1] D. Stover, H.P. Buchkremer, S. Uhlenbruck, Ceram. Int. 30 (2004) 1107. [2] B.C.H. Steele, J. Mater. Sci. 36 (2001) 1053. [3] J.M. Ralph, C. Rossignol, R. Kumar, J. Electrochem. Soc. 150 (2003) A1518. [4] S. Simner, J.F. Bonnett, N.L. Canfield, K.D. Meinhardt, V.L. Sprenkle, J.W. Stevenson, J. Power Sources 113 (2003) 1. [5] S. Simner, J.F. Bonnett, N.L. Canfield, K.D. Meinhardt, V.L. Sprenkle, J.W. Stevenson, Electrochem. Solid-State Lett. 5 (7) (2002) A173. [6] G.Ch. Kostogloudis, Ch. Ftikos, Solid State Ion. 126 (1999) 143. [7] A. Esquirol, N.P. Brandon, J.A. Kilner, M. Mogensen, J. Electrochem. Soc. 151 (2004) A1847. [8] A. Mai, V.A.C. Haanappel, S. Uhlenbruck, F. Tietz, D. Stover, Solid State Ion. 176 (2005) 1341. [9] J.B. Goodenough, in: C.N.R. Rao (Ed.), Solid State Chemistry, Dekker, NY, 1974, p. 215. [10] R.J.H. Voorhoeve, in: J.J. Burton, R.L. Garton (Eds.), Advanced Materials in Catalysis, Academic Press, NY, 1977, p. 129. [11] S.E. Dann, D.B. Currie, M.T. Weller, J. Solid State Chem. 109 (1994) 134. [12] R.J. Bell, G.J. Millar, J. Drennan, Solid State Ion. 131 (2000) 211. [13] Y. Teraoka, M. Yoshimatsu, N. Yamazoe, T. Seiyama, Chem. Lett. (1984) 893. [14] F. Zheng, L. Pederson, J. Electrochem. Soc. 146 (1999) 2810. [15] S. Otoshi, H. Sasaki, H. Ohnishi, M. Hase, K. Ishimaru, M. Ippommatsu, T. Higuchi, M. Miyayama, H. Yanagida, J. Electrochem. Soc. 138 (1991) 1519.

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