Characterization of oxygen-deficient SrCoO3−δ by electron energy-loss spectroscopy and Z-contrast imaging

Solid State Ionics 130 (2000) 71–80 www.elsevier.com / locate / ssi

Characterization of oxygen-deficient SrCoO 32d by electron energy-loss spectroscopy and Z-contrast imaging a, b ,1 c b S. Stemmer *, A. Sane , N.D. Browning , T.J. Mazanec a

Department of Mechanical Engineering and Materials Science, Rice University, MS 321, 6100 Main Street, Houston, TX 77005 -1892, USA b BP Chemicals /Amoco, Naperville, IL 60566, USA c Physics Department, University of Illinois at Chicago, Chicago, IL 60607 -7059, USA Received 27 September 1999; received in revised form 5 November 1999; accepted 22 November 1999

Abstract Atomic resolution Z-contrast imaging and electron energy-loss spectroscopy (EELS) were used to study the defect structure of oxygen deficient strontium cobaltite (SrCoO 32d ) after repeated oxidation / reduction cycles at elevated temperatures. Ordered microdomains were found in the reduced sample. Z-contrast imaging showed ordering in alternate CoO 2 planes, characteristic for a brownmillerite-type vacancy ordered structure. EELS showed that the average valence state of cobalt had decreased after testing. In the tested sample, cobalt was found in 1 3 and 1 2 valence states, respectively. In addition, cobalt oxide precipitates were found at grain boundaries and in triple pockets in this sample, indicating a cobalt deficiency of the grains. The mechanisms for the formation of the observed microstructure are discussed.  2000 Elsevier Science B.V. All rights reserved. Keywords: Brown millerite structure; Oxygen deficient strontium cobaltite; Ceramic membranes

1. Introduction Perovskite-type oxides that combine a high electronic and ionic conductivity are very promising materials for use as dense ceramic membranes for oxygen separation [1]. Many of the physical properties of transition metal perovskite-type oxides (ABO 3 ), including the oxygen transport properties, *Corresponding author. Tel.: 1 1-713-348-3546; fax: 1 1-713348-5423. E-mail address: [email protected] (S. Stemmer) 1 Now at: Advanced Ceramics Corporation, Cleveland, OH, USA.

are related to the extent of oxygen nonstoichiometry [2], which depends on temperature and oxygen partial pressure. The oxygen nonstoichiometry can also be controlled by the substitution of lower valence cations on the A-site (e.g. of Sr 12 in place of La 13 ). In many oxygen deficient perovskite-type oxides, superstructures of ordered vacancies, such as the well-characterized brownmillerite structure shown in Fig. 1, are formed [3]. At elevated temperatures the oxygen vacancies are believed to have a tendency to disorder. Such order–disorder transitions can influence the oxygen permeability [4]. Another important issue in the development of perovskite-type oxygen conducting membranes is the

0167-2738 / 00 / $ – see front matter  2000 Elsevier Science B.V. All rights reserved. PII: S0167-2738( 99 )00309-4

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Fig. 1. Vacancy ordering in the brownmillerite structure in alternate (001) BO 2 planes along the [110] direction of the cubic perovskite parent structure, after Ref. [3]. The A-site cations, located at the corners, are not shown.

structural, chemical and mechanical stability of the materials at high temperatures and in reducing environments. To solve these problems, it is important to understand the factors that determine the microstructure of these materials under the conditions of a given application. The goal of the present paper is to investigate the capabilities of a combination of transmission electron microscopy spectroscopy and imaging techniques for this purpose, using a model material, SrCoO 32d , subjected to reduction cycles at high temperatures, as an example. SrCoO 32d can be highly oxygen deficient, and various phases with different crystal structure, composition and electronic states of cobalt exist [5–8]. Its physical properties, and the fact that SrCoO 32d is also the parent phase of other potential candidates for oxygen conducting materials, for example

SrCo 12x B x9 O 32d (B9 5 Cr, Fe, Co and Cu) [4] make studies of defects in this material very interesting. SrCoO 3 is not stable in air at elevated temperatures and loses oxygen to approximately SrCoO 2.5 , which during slow cooling in air oxidizes to SrCoO 2.51d [9]. SrCoO 2.5 is known to assume two crystallographic structures, a brownmillerite-type microstructure at high temperatures, and a ‘lowtemperature hexagonal’ form (H-SrCoO 32d ) [10]. The crystal structure and magnetic properties of H-SrCoO 32d were the subject of several investigations [11–13]. Harrison et al. [13] resolved the discussion by showing that during cooling of SrCoO 2.5 in air, phase separation in Sr 6 Co 5 O 15 and Co 3 O 4 occurs, and the authors were able to determine the structure of Sr 6 Co 5 O 15 (‘H-SrCoO 32d ’). Recently, Vashook et al. investigated phase transitions of SrCoO 2.52d from a rhombohedral (d # 0.16) through a cubic ‘pseudo-perovskite’ and eventually to a perovskite structure (for d 5 $ 0.21), as a function of oxygen partial pressure and temperature [6]. The authors suggest that the cubic ‘pseudoperovskite’ (0.16 # d # 0.21) contains ordered microdomains of a lower symmetry phase. Electron energy-loss spectroscopy (EELS) in transmission electron microscopy (TEM) can be used to analyze unoccupied states above the Fermi level and to determine the valence states of transition metal atoms [14,15]. High-resolution imaging techniques, such as conventional high-resolution TEM (HRTEM) and Z-contrast in scanning TEM (STEM), can provide information about ordered superstructures. In this paper, we use these techniques to investigate the structural stability of an oxygen deficient SrCoO 32d ceramic that has been cycled between oxidizing and reducing atmospheres at high temperatures.

2. Experimental Stoichiometric amounts of Sr acetate and Co 3 O 4 were ground and calcined at 8008C for 6 h, followed by grinding and sintering of the compressed powder at 12008C for 1 h in air. Electron diffraction showed that after cooling the sample had H-SrCoO 32d structure, more correctly described as a mixture of Co 3 O 4 and Sr 6 Co 5 O 15 [13] (in the following referred to as

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‘as-sintered’ sample). In order to test the structural stability of the material at high temperatures and in reducing environments, the following experiment was performed. The sample was heated up in air at a rate of 5 K / min. During heating, a phase transformation, most likely from the ‘H-SrCoO 32d ’ structure to a cubic or pseudo-perovskite structure, was observed at about 7508C. After reaching 10148C, the atmosphere was changed three times from air to nitrogen to air, followed by cycling twice between air and CO 2 . The sample expanded in N 2 or CO 2 due to the creation of oxygen vacancies, and contracted in air due to oxygen sorption [16]. Before each change in environmental conditions the sample was allowed to achieve equilibrium (no expansion or contraction). After the last step (CO 2 ) the sample was cooled in N 2 at a rate of about 5 K / min. The relative expansion / contraction were 0.17760.007% for air / nitrogen and 0.12860.0044% for air / CO 2 [17]. The weight loss of the sample after testing was 1.68% [17]. If the weight loss is attributed to oxygen loss, and the starting composition of the sample was approximately SrCoO 2.51d , the sample had a composition of SrCoO 2.311d after testing. Details and theoretical background of this experiment will be reported elsewhere [17]. The present paper will concentrate on the microstructural characterization of the tested sample. To this purpose, TEM specimens were prepared before and after testing, using standard preparation methods, with argon ion-milling as the final step. TEM investigations were performed using a 200 kV transmission electron microscope (JEOL JEM 2010F) equipped with a field-emission gun, an annular dark-field detector and a post-column imaging filter (Gatan GIF200). This microscope is capable of achieving sub-0.2 nm probe sizes [18] for microanalysis and incoherent Z-contrast lattice imaging. Z-contrast images are more readily interpretable than conventional HRTEM images, e.g. the bright intensity maxima correspond to the position of the atom columns [19]. The angular range collected by the annular detector was about 35 mrad. EELS was employed for studying differences in bonding and valence states of cobalt. Valence state changes of transition metals with unoccupied 3d states can be analyzed by measuring the intensity

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ratios of the L 3 and L 2 lines, which are sensitive to the number of unoccupied d-states [15]. Here, cobalt L 3 / L 2 intensity ratios were determined by two different methods. The first method measures the area underneath the positive peaks in second difference spectra [20] (denoted as derivative method). The second method measures the area underneath the peaks of spectra with the continuum part of atomic cross-section distributions subtracted, as suggested by Kurata and Colliex [21] (denoted as subtraction method). Details of both quantification procedures are described in Ref. [22]. Oxygen K near-edge fine structure (ELNES) were also recorded for additional information about changes in the local atomic environment, as the ELNES of this edge can be interpreted to reflect the density of unoccupied states above the Fermi level [14].

3. Results After testing, SrCoO 32d showed two prominent microstructural features: cobalt oxide precipitates at grain boundaries and in triple junctions, and ordered microdomains. Fig. 2 shows such a precipitate in a

Fig. 2. Low-magnification TEM micrograph of the sample after testing, showing a cobalt oxide precipitate at a triple junction. Microcracks along grain boundaries can also be observed.

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triple junction. Energy dispersive X-ray analysis (EDS) showed that these precipitates are rich in cobalt and oxygen, and contain no detectable amount of strontium. Fig. 3(a) shows a HRTEM image of a grain oriented along the [010] p axis (p denotes the cubic perovskite reference lattice), and Fig. 3(b) a selected area diffraction (SAD) pattern of this region. The insets in the HRTEM image are Fourier transforms (FFTs) of two domains, and exhibit a doubling of the perovskite unit cell along [001] p , also visible in the SAD pattern. In the electron diffraction pattern, the superlattice reflections originate from two sets of domains and are indicated by

arrows in Fig. 3(b). The domains are rotated about 908 with respect to each other. Fig. 4(a) shows an incoherent Z-contrast lattice image of one domain, recorded along [010] p . Fig. 4(b) is the same image after Fourier filtering to remove the noise in the image. Under the experimental conditions, the image contrast is sensitive to the atomic number, and strontium columns can be distinguished from columns containing alternate cobalt and oxygen atoms (the inset in Fig. 4(b) shows the unit cell in this projection). The image intensity in every second CoO 2 plane (indicated by arrows) is decreased significantly. This change in

Fig. 3. (a) HIRTEM image of a grain in the tested sample along a [010] p axis; the insets in the HRTEM image are Fourier transforms (FFTs) taken from two domains; (b) a selected area diffraction pattern of this region with superlattice reflections originating from two sets of domains (a and b) as indicated by arrows.

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Fig. 4. (a) Z-contrast lattice image of one domain, recorded along [010] p ; (b) is the same image after Fourier filtering. The inset in (b) shows a possible model for the unit cell in this projection.

Fig. 5. (a) Z-contrast lattice image of one domain, recorded along k110l p ; (b) is the same image after Fourier filtering with an inset showing a possible model of the unit cell in this projection.

contrast could be due to missing atom columns in these CoO 2 planes, e.g. oxygen. However, oxygen is known to contribute only very little to the image contrast [19] in such images, and other, or additional, structural origins for this particular change in the image contrast should be considered. For example, a shift of the cobalt atoms to achieve a tetrahedral coordination in planes with missing oxygen [3] might also contribute to a change in image contrast. Another possibility might be missing cobalt. Without image simulations, the relative contribution of each of these possibilities is difficult to quantify. Fig. 5(a) shows a Z-contrast image of a grain along a [110] p zone axis, Fig. 5(b) is the same image after Fourier filtering. The resolution in this image is sufficient to resolve the cobalt and strontium columns along the [001] p direction. However, the spacing between the oxygen columns and the cobalt columns (about 0.14 nm) in this projection is smaller than the spatial resolution (probe size) in this image. As in the [010] p projection, a decrease in image

intensity can be observed in alternate CoO 2 planes (indicated by arrows). Fig. 6(a) shows Co L-edges recorded from the as sintered and from the tested sample. Fig. 6(b) shows the L 3 / L 2 intensity ratio obtained by the derivative method for all experiments performed. The L 3 / L 2 intensity ratio is higher in the tested sample. Average cobalt L 3 / L 2 ratios of the two samples are summarized in Table 1. Although the values obtained using the derivative method are somewhat different from those obtained with the subtraction method, both clearly show an increase in L 3 / L 2 intensity ratio after testing. In literature, L 3 / L 2 ratios have been measured for several cobalt oxides [23,24]. Generally, the L 3 / L 2 ratio increases with decreasing valence state of Co. Comparison with L 3 / L 2 ratios measured by Wang and Yin [23] for different cobalt oxides renders a cobalt valence state of about 1 2.9 (60.3) for the as sintered H-SrCoO 32d , and a valence state of 1 2.4 (60.3) for the tested SrCoO 32d . The errors given in parenthesis have been estimated from the

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Fig. 6. (a) EELS of cobalt L-edges of SrCoO 32d before (as sintered), and after treatment in oxidizing / reducing atmospheres at 10148C (tested). The spectra were recorded with an energy dispersion of 0. 1 eV/ channel, and an energy resolution at the zero-loss peak of no more than 1.2 eV (full width at half maximum). The acquisition time was 20 s; the nominal probe size was 1.5 nm. (b) Cobalt L 3 / L 2 intensity ratios for the tested and untested sample obtained with the derivative method.

Table 1 Cobalt L 3 / L 2 ratios and statistical error (in parenthesis) obtained from the as sintered and the tested sample. L 3 / L 2 ratios were obtained by the two methods described in Ref. [22]. The difference in the L 3 / L 2 ratios obtained by the two methods can serve as an estimate of the systematic error in determining the L 3 / L 2 ratio by these techniques. The experimental conditions are given in Fig. 6 Sample

L 3 / L 2 (derivative method)

L 3 / L 2 (subtraction method)

SrCoO 32d (as sintered) SrCoO 32d (tested)

2.87 (60.26) 3.86 (60.11)

3.48 (60.35) 4.35 (60.26)

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different L 3 / L 2 ratios obtained by the two quantification methods used to determine the L 3 / L 2 ratio (this error is larger than the statistical errors given in Table 1). Other factors besides the valence state can influence the L 3 / L 2 ratio, especially the spin states [25]. Brownmillerite-type SrCoO 2.5 contains highspin Co 31 in both the octahedral and the tetrahedral sites, whereas the hexagonal phase contains low-spin Co 31 on the octahedral sides [9]. As suggested by Botton et al. [22], Fig. 6(a) shows the normalized (with respect to the continuum after the edge) cobalt L 3,2 edges of both samples. This plot shows that the integrated white line intensity is different before and after testing, indicating a change in unoccupied 3d states at the cobalt sites and a change in valence state. At least in our case, the change in spin states seems to have little influence on the Co L 3 / L 2 ratios. This was also observed by Wang and Yin [23]. Fig. 7 shows oxygen K-edges obtained from the as-sintered and the tested sample. The spectra were acquired using a nominal probe size of 1.5 nm, and an energy dispersion of 0.1 eV/ channel. The two spectra show, among other differences in the fine structure, a difference in the relative height of the pre-peak (at around 528 eV, denoted a in Fig. 7) with respect to the first peak (denoted b). The pre-peak in

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oxygen K-edges in 3d transition metal oxides is known to arise from hybridization of oxygen p states with the transition metal d density of states and is correlated with the 3d-band occupancy of the transition metal [26]. The decrease in height of the prepeak in the tested sample confirms the decrease in valence state also observed from the cobalt L-edges. A second set of EELS experiments was performed using a small probe size of 0.2 nm, and a larger energy dispersion of 0.5 eV (to record both Co L-edge and O K-edge in a single spectrum). In these experiments, an attempt was made to position the probe in a k010l Z-contrast image at different positions within one unit cell. During most experiments the sample drifted during the acquisition time for one spectrum (10 s). However, a number of spectra showed a significantly higher or lower than average cobalt L 3 / L 2 ratio, respectively. These spectra were selected for further analysis. In a first approximation, we assume that these spectra indicate that different valence states exist within one unit cell. To further investigate this hypothesis, two EELS spectra are shown in Fig. 8, each being a sum of six spectra with cobalt L 3 / L 2 ratios higher than average (low valence state) or a lower than average (high valence state), respectively. The spectra that showed

Fig. 7. EELS of oxygen K-edges of as sintered and tested SrCoO 32d . The spectra were recorded with an energy dispersion of 0.1 eV/ channel, and an energy resolution at the zero-loss peak of no more than 1.2 eV (full width at half maximum). The acquisition time was 20 s, the nominal probe size was 1.5 nm.

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Fig. 8. EELS spectra that gave a high L 3 / L 2 ratio (low valence state) or a low L 3 / L 2 ratio (high valence state) when using a probe size of 0.2 nm. Each spectrum is the sum of six spectra, acquired with an exposure time of 10 s.

a high valence state of Co (as derived from the L 3 / L 2 ratios) also show a shift of the Co L-edges towards higher energy losses. Such energy shifts are characteristic for higher valence states [27]. The Co valence states derived from the L 3 / L 2 intensity ratios for these spectra are approximately 3 1 and 2 1 , respectively. In addition, the oxygen K edges of the two spectra differ in the relative height of the pre-peak and in other features of their ELNES. The observed energy shift of the cobalt L-edges and the correlated changes in the oxygen K-edge ELNES provide additional support that the measured L 3 / L, intensity ratios reflect the existence of different cobalt valence states within the domains.

4. Discussion Measurements of cobalt L 3 / L 3 intensity ratios and correlated changes in the fine-structure of oxygen K-edges in EELS spectra of SrCoO 32d show that after repeated oxidation / reduction cycles at approximately 10008C the average cobalt valence state has decreased from an initial value of about 1 3 to about 1 2.5. Within the experimental error, the change in cobalt valence state is consistent with the change in the oxygen nonstoichiometry that was estimated from the weight loss of the sample, from SrCoO 2.51d , before testing to SrCoO 2.311d after testing, assuming the valences are 1 2 for Sr and 2 2 for O. Further investigations of the oxygen

K-edge ELNES are underway to interpret the changes in the fine structure for changes in the cobalt coordination. The reduced cobalt valence state reflects the oxygen nonstoichiometry of the material in equilibrium with CO 2 after several oxidation / reduction cycles at the testing temperature. The result is consistent with oxygen desorption studies of Teraoka et al. [28] that showed a so-called b desorption peak at high temperatures in SrCoO 2.5 and other perovskite-type materials, correlated with a reduction of the transition metal ion from B 31 to B 21 . The tested sample in our study clearly shows the presence of ordered microdomains at room temperature. Brownmillerite-type microdomains have been reported in other materials [29,30]. The formation of microdomains has also been used to explain anomalies in the oxygen permeation behavior of SrCoO 32d based materials [4]. In the SrCoO 32d system, microdomains have not been previously imaged by HRTEM. However, Takeda et al. [5] suggested that the apparent cubic symmetry of their SrCoO 2.29 in X-ray diffraction might be caused by microdomains, and Vashook et al. [6] suggest the same for their ‘pseudo-perovskite’ in SrCoO 2.52d (0.16 # d # 0.21). The Z-contrast images in our study indicate missing atom columns and / or structural rearrangements in alternate CoO 2 planes; both are characteristic for brownmillerite type vacancy ordered structures. Although other authors have observed that the solid solution range of the brownmillerite phase in this system extends into the

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Co 21 / Co 31 mixture region [5], the average cobalt valence state measured for this sample ( 1 2.5) is much lower than in an ideal brownmillerite SrCoO 2.5 . At this nonstoichiometry, the reduction of the sample has probably reached the stability limit of the material as large cobalt oxide precipitates are observed at the grain boundaries and in triple junctions after testing. Cobalt deficiency of the grains might be a possible reason why the vacancy ordered, perovskite derived structure is still observed. It should be noted that, based on the observations of other authors [13], it is expected that some degree of cobalt deficiency and a source for divalent Co exists already in the untested sample, and that this deficiency has further increased by the cyclic treatment. However, without image simulations, it is not possible to judge whether a cobalt deficiency in alternate CoO 2 planes might also be the reason for the strong contrast observed in the Z-contrast images. The derivation from the ideal brownmillerite composition also raises the question why microdomains are formed in our sample. In literature, the formation of microdomains is proposed to serve as a mechanism to accommodate oxygen excess (relative to the ideal stoichiometry ABO 2.5 ) by incorporation in the domain walls [29,30]. Accommodation of oxygen excess cannot be the cause for the microdomains at the nonstoichiometry measured in our sample. Instead, likely reasons include strain relaxation during transformation from a disordered cubic perovskite (to preserve an average cubic symmetry), transformation kinetics, or the large derivation from the ideal nonstoichiometry. To distinguish between these possibilities, in situ experiments should be performed to determine whether the microdomains are present at high temperatures.

5. Conclusions In summary, we have used a combination of spectroscopy and imaging techniques in a scanning transmission electron microscope to elucidate the effects of repeated reduction / oxidation cycles on the structure of an oxygen deficient perovskite-type SrCoO 32d ceramic. Our results show that decomposition, microdomain formation, reduction of cobalt from a 1 3 to a 1 2 valence state, and cobalt as well

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as oxygen nonstoichiometry are important microstructural aspects of this material under application conditions.

Acknowledgements The authors would like to acknowledge BP Chemicals /Amoco for support of this work. The work at BP Chemicals was supported under a NIST ATP Cooperative Agreement 70NANBSH106S. We also acknowledge the use of the microscopy facilities at the RRC at the University of Illinois at Chicago (NSF contract DMR-9601792). S.S. would like to thank Dr. Yan Xin for useful discussions.

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