Relationship between cation substitution and stability of perovskite structure in SrCoO3–δ-based mixed conductors

Solid State Ionics 177 (2007) 3433 – 3444 www.elsevier.com/locate/ssi

Relationship between cation substitution and stability of perovskite structure in SrCoO3–δ-based mixed conductors T. Nagai ⁎, W. Ito, T. Sakon Advanced Technology Research Laboratories, Nippon Steel Corporation, Futtsu, Chiba 293-8511, Japan Received 6 July 2006; received in revised form 29 September 2006; accepted 26 October 2006

Abstract In SrCoO3–δ (SC)-based mixed conductors, cation substitution is necessary to stabilize the high-oxygen permeable perovskite structure, but the dose of the substitution should be minimized to avoid degradation of permeability. To clarify the relation between the substitutional cation and the perovskite stability of the mixed conductor, SC-based oxides, (La0.1Sr0.9)CoO3–δ and Sr(Co0.9X0.1)O3–δ where X was Ni, Cu, Zn, Cr, Fe, Al, Ga, In, Ce, Ti, Zr, Sn, V and Nb, were studied. The appearance of the low oxygen permeable 2H–BaNiO3-type SrCoO2.52 phase during the preparation of the powder samples or post-annealing in oxygen was investigated and the tendency to transform to the hexagonal phase was evaluated. The sequence of the perovskite stability upon the substituting cation for SC was Ni; Cu; Zn; In; CebCr; Al; Ga; Zr; Sn; VbLabFebTibNb: Thermogravimetry revealed that a rise in the valence of the substitutional cation increases the oxygen content, 3–δ, and enhances the stability of the perovskite structure. Oxygen permeability of the ceramic disk samples of SC, (La0.1Sr0.9)CoO3–δ (LaSC), Sr(Co0.9Fe0.1)O3–δ(SCFe), Sr(Co0.9Ti0.1)O3–δ (SCTi), Sr(Co0.9Nb0.1)O3–δ (SCNb) and Sr(Co0.8Fe0.2)O3–δ (SCF1082) were measured: the order of oxygen permeability at 900 °C was SCNb≥SCNSCTiNSCFeNLaSCNSCF1082: Comparison of the sequence of the perovskite stability and oxygen permeability suggests that oxygen permeability is not in a trade-off relation against the stability of the perovskite phase. Within the measured samples, SCNb exhibited the highest stability of the perovskite structure, the highest oxygen permeability of 4.24 cm3/min/cm2 at 900 °C, good gas-tightness and steady thermal expansion of 22.8 ppm/°C. Niobium substitution on the Co-site was found for the first time to be effective in preparing an SrCoO3–δ based mixed conductor for application as an oxygen separation membrane. © 2006 Elsevier B.V. All rights reserved. Keywords: Perovskite; Phase transition; Stability; Oxygen permeability; Membrane

1. Introduction Mixed electronic/oxide-ionic conductors have been widely studied as membrane materials for oxygen separation. Such membranes can be used in pressure driven oxygen generators and membrane reactors for partial oxidation, e.g., syngas production from methane [1]. The efficiency of the membrane ⁎ Corresponding author. Tel.: +81 439 80 3065; fax: +81 439 80 2746. E-mail address: [email protected] (T. Nagai). 0167-2738/$ - see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2006.10.022

electrochemical systems depends on the oxygen permeability of the utilized mixed conductors and thus the development of mixed conductors with high oxygen permeability is the key to realizing oxygen separation systems. Teraoka et al. found that the rates of oxygen permeation through perovskite-type ceramic membranes, SrCo0.8Fe0.2O3–δ and La0.6Sr0.4CoO3–δ, are one or two orders of magnitude greater than the rate of stabilized zirconia under short-circuited conditions. Since then, mixed-conductive perovskite-type oxides have been thoroughly investigated as membrane materials [2], and a

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number of perovskite related mixed conductors have been examined [2–17]. According to the literature, mixed conductors with high oxygen permeability will be found among perovskites containing cobalt, such as SrCoO3–δ (SC)-based solid solutions [1,18]. One problem with cobalt containing mixed conductors is the structural stability under a low oxygen activity environment. For example, it is difficult to apply SC-based mixed conductors for application as membrane reactors, because one side of the membrane is exposed to a reductive gas such as CH4. On the other hand, oxygen activity becomes comparatively high in oxygen generators. In the generators the membrane separates the oxygen from the pressurized air, for example 20.4 atm: the oxygen partial pressure of each side of the membrane surface, Po2, is 4.2 and 1 atm [19]. In these systems, the driving force of the oxygen permeation, log(Po2high/Po2low), is extremely lower than in the case of membrane reactors. Thus the membrane materials applied for oxygen generators need to have higheroxygen-permeability. Consequently, it is reasonable to exploit the membrane material for oxygen generators among a group of SC-based mixed conductors. As reported for (La, Sr)(Co, Fe)O3–δ, the oxygen permeability of SC-based solid solutions increases with decreasing the concentration of the substituting cations for Sr and Co, however it peaks and begins to decrease in the vicinity of the ultimate composition, SC [2]. This is because SC is changed to a hexagonal phase with the 2H–BaNiO3-type structure below about 900 °C [3,5]. The hexagonal phase has oxygen vacancy ordering and it is almost non-oxygen permeable [5,20]. These results suggest the importance of controlling the phase transition to the BaNiO3-type structure. The high-oxygenpermeable perovskite structure must be stabilized by substituting SC with proper cations to form a solid solution. Furthermore, the dose of the substitutional cation should be minimized to avoid degradation of the oxygen permeability. In this study, we focused on the change in the stability of perovskite upon the substitutional cation for an SC-based mixed conductor. SC-based mixed conductors have been widely studied. Systems such as (La, Sr)(Co, Fe)O3–δ including Sr(Co, Fe)O3–δ and (La, Sr)CoO3–δ, (Ba, Sr)(Co, Fe)O3–δ, Sr(Co, Cu)O3–δ, Sr(Co, Ni)O3–δ, Sr(Co, Cr)O3–δ, Sr(Co, Ti)O3–δ, Sr(Co, Mn)O3–δ, Sr(Co, Fe, Sn)O3–δ, Sr(Co, Fe, Ge)O3–δ, Sr(Co, Fe, Zr)O3–δ, Sr(Co, Ce)O3–δ, and Sr(Co, V)O3–δ have been investigated [2,3,5–7,9–17]. However, comparative information on the relation between the change in the stability and the kinds of substitutional cation is lacking. Such information is valuable for elaborating SC-based mixed conductors. By using a cation having a higher effect on stabilizing the perovskite structure, the dose of the substitution for SC can be minimized, which will lead to realizing both high permeability of oxygen and high stability of perovskite. In the present study, the dose of the cation substitution was fixed at 10 mol% of the A- or B- site, and the phase compositions of the samples were investigated in order to clarify the change in the stability of the perovskite structure upon the substituting cation. Oxygen nonstoichiometry, 3–δ, was measured and discussed in connection with the mechanism of

the perovskite stabilization upon the cation substitution for SC. For some of the SC-based oxides, oxygen permeation and thermal expansion were measured. Besides the investigation, we hypothesized that a substitutional cation with a higher valence would be more effective in stabilizing the perovskite of an SC-based mixed conductor. For the A-site, i.e. the Sr2+-site, only trivalent cations such as rare earth elements are candidates for the higher-valence substitutional cation. For the B-site, i.e. the Coz+-site where z b 4, four- and fivevalent cations are candidates. Substitution by four-valent cations such as Ti4+, Sn4+ and Zr4+ has been investigated, and substitution by a five-valent cation has been reported only for the Sr(Co, V5+ )O3–δ system by Kharton et al.: Sr(Co, V 5+ )O3–δ contains two phases and its oxygen permeability is higher than that of SC. However, the stability of the perovskite of this system is unclear [7]. Among the studied solid solutions, Sr(Co0.9Nb0.1)O3–δ exhibited the most stable perovskite structure and the highest oxygen permeability. Niobium substitution on the B-site was found to be effective in preparing an SrCoO3–δ based mixed conductor for oxygen separation [21]. 2. Experimental The specimens were prepared according to a solid reaction and sintering process. The synthesized compositions were SrCoO3–δ, (La0.1Sr0.9)CoO3–δ and Sr(Co0.9X0.1)O3–δ where X was Ni, Cu, Zn, Cr, Fe, Al, Ga, In, Ce, Ti, Zr, Sn, V and Nb. In the following, each composition is abbreviated to SC, LaSC, SCNi, SCCu, SCZn, SCCr, SCFe, SCAl, SCGa, SCIn, SCCe, SCTi, SCZr, SCSn, SCV and SCNb, respectively. Reagentgrade SrCO3, Co3O4, NiO, CuO, ZnO, M2O3 (M = La, Cr, Fe, Al, Ga and In), CeO2, TiO2, ZrO2, SnO2, V2O5 and Nb2O5 were weighed and mixed by ball milling for 3 h with ZrO2 balls. Ethyl alcohol was used as the medium. The powder mixtures were calcined at 900 °C for 12 h in air and finely ground by ball milling. Then each powder was sintered at 1200 °C or 1150 °C for 5 h in air. After the sintering, the powder was ground again. Half of the powder was annealed at 850 °C for 24 h in oxygen gas flow. The heating and cooling rate of the sample in the sintering and the annealing was 200 °C/h. X-ray powder diffraction analysis of the powder samples was carried out at room temperature using a RIGAKU RAD-3C diffractometer. Cu–Kα radiation was used as the X-ray source. Oxygen nonstoichiometry, 3–δ of the powder samples were measured by thermogravimetry using a RIGAKU Thermo-plus TG8120 thermogravimetric and differential thermal analyzer (TG-DTA). The gas flow was controlled to 100 cm3/min by mass flow controllers STEC SEC-400 MARK3. First, 3–δ in the “initial condition”, 3–δ0, was measured: a fresh sample of about 40 mg was annealed at 1000 °C for 0.5 h in air flow and cooled to 80 °C in the TG-DTA; next, the atmosphere was changed to 4% hydrogen/96% argon mixture gas and the sample was reduced by heating to 1100 °C. By heat treatment the SCbased oxides were decomposed and Co, Fe and Cu were reduced to metal, while other cations were not reduced under the condition: therefore the observed mass loss is apportioned to

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the change in oxygen content upon the reduction of Co, Fe and Cu from the initial condition to metal. With the obtained value of 3–δ0, 3–δ under different conditions could be determined by the weight change from the initial condition. 3–δ in oxygen flow was measured in the present study. A fresh sample was annealed at 1000 °C for 0.5 h in air flow and cooled to 80 °C in the TG-DTA to reach the initial condition; then the atmosphere was changed to oxygen and the sample was heated to 1000 °C, maintained for 0.5 h and cooled to 100 °C. 3–δ in oxygen was determined from the weight measured during the cooling step. The heating and cooling rate was 3 °C/min. For SC, LaSC, SCFe, SCTi and SCNb, ceramic samples were also prepared. The calcined and ball milled powders were isostatically pressed into pellets. The pellets were sintered at 1150 °C, 1200 °C and 1250 °C for 5 h in air. The densities of the sintered pellets were measured according to Archimedes' method. The ceramic sample whose density was the highest of the three was sliced and polished into discs of approximately 1 mm in thickness. The discs were used for oxygen permeability measurement. The oxygen permeation was measured using the apparatus shown in Fig. 1. The disk-shaped sample was placed between two concentric Al2O3 tubes. The sealing between the outer Al2O3 tubes and the sample was achieved by glass rings. The side seal of the sample was also achieved by a glass seal. Prior to the measurement, the test section was heated to 700 °C to form the glass seals. After sealing, it was heated from 750 °C to 900 °C or 950 °C. The temperature was held at interval of 50 °C and the oxygen permeability was measured. The permeability measurements were conducted in both heating and cooling processes.

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During the permeability measurement, the upper side of the sample was supplied with a mixture of oxygen and helium as the feeding gas. The gas flow was controlled by mass flow controllers LINTEC MC-1100. The oxygen partial pressure of the feeding gas was kept at 0.21 atm, 0.5 atm and 1 atm in order to evaluate the dependence of permeability on (PO2)1/n. The total gas flow rate of the upper side was kept at 40 cm3/min. The lower side of the sample was supplied with argon as the sweep gas. The oxygen potential of the outlet gas was monitored by oxygen analyzer Toray LC-800. The argon gas flow rate was controlled at 40 cm3/min by a mass flow controller. The sample was maintained under fixed conditions for more than 1 h until the oxygen potential of the outlet gas reached equilibrium. The permeation of oxygen through the sample was evaluated from the gas flow rate and the oxygen potential of the outlet gas. The leak rate through the sample was estimated from the amount of helium in the outlet gas, which was measured using a quadrapole gas analyzer ANELVA AQA-100 MPX. The oxygen permeation was corrected by subtracting the amount of oxygen due to the leak. The efficiency of the side seal was confirmed by the measurement of the nitrogen in the outlet gas. The oxygen permeability was compared under the same conditions as follows: the sample thickness was 1 mm; oxygen partial pressure of the feeding and sweeping gas was 0.21 atm and 0.01 atm. These values were obtained under the assumption that the oxygen permeation of the samples is in inverse proportion to the sample thickness, and it depends on the oxygen partial pressure by the power law, (PO2–high)1/n–(PO2–low)1/n. Dense ceramic bars of 5 mm and 5 mm and 20 mm were also prepared for LaSC, SCFe, SCTi and SCNb and submitted to thermo-dilatometric measurement using RIGAKU Thermo Plus 2 TMA8310. An Al2O3 bar was used as a reference for the TMA measurement. After heating and keeping at 1000 °C for 30 min, the contraction curve was measured during cooling to 50 °C. The cooling rate was 1 °C/min. The sample environment was controlled by flowing air or oxygen at 100 cm3/min. 3. Results and discussion

Fig. 1. Schematic drawing of the oxygen permeability measurement apparatus.

X-ray diffraction patterns of SC, LaSC, SCNi, SCCu, SCZn, SCCr, SCFe, SCAl, SCGa, SCIn, SCCe, SCTi, SCZr, SCSn, SCV and SCNb were measured for the samples sintered at 1200 °C for 5 h in air. Fig. 2 shows the pattern of SC, SCNi, SCCu, SCZn, SCIn and SCCe. The XRD patterns in the figure show that all of the six samples contain SrCoO2.52, JCPDS 401018, BaNiO3-type hexagonal phase [22]. In Fig. 2, the peaks corresponding to the hexagonal phase are marked with “h”. The XRD patterns of SC and SCCu are in good agreement with SrCoO2.52 and these samples contain only the hexagonal phase. The indices of the diffraction peaks on the basis of the SrCoO2.52 hexagonal are shown above the peaks of SC after the JCPDS card. Since some weaker peaks are left unidentified in the card, some peaks are marked with “h” but not indexed. Though the crystal structure of SC is still under discussion [23– 25], it will be referred as SrCoO2.52 hexagonal in the following. The XRD patterns of SCNi, SCZn, SCIn and SCCe show that these samples contain both the hexagonal phase

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The subscript “c” and “o” denotes the cubic perovskite and the orthorhombic brownmillerite, respectively. For LaSC, SCCr, SCTi, SCZr, SCSn, SCV and SCNb, the main feature of the XRD patterns can be explained by the cubic phase with perovskite structure. The indices of the diffraction peaks on the basis of the cubic perovskite are shown above the peaks of LaSC. Additionally, some small peaks are present in the

Fig. 2. X-ray diffraction patterns after sintering at 1200 °C for 5 h in air for SC, SCNi, SCCu,SCZn,SCIn and SCCe. Peaks shown by “h” correspond to SrCoO2.52; “N” to Sr2Ni2O5; “z” to ZnO;“I”to SrIn2O4; “e” to SrCeO3. Inscriptions above the peaks for SC are the Miler's indexes (hkl) of the hexagonal SrCoO2.52.

and subsidiary phase. SrCo1–xNixO3–δ where 0.1 ≤ x ≤ 0.2 has been reported to form a solid solution of hexagonal structure [6]. In the present study, however, traces that could be assigned to Sr2Ni2O5 (JCPDS 28-1242) appear in the pattern of SCNi in Fig. 2. In addition, some peaks that agree with the perovskite structure also appear in the diffraction pattern and are marked with “p” in the figure. In the pattern of SCZn, SCIn and SCCe, additional peaks corresponding to ZnO (JCPDS 361451), SrIn2O4 (33-1336) and SrCeO3 (36-980) are observed, respectively. In contrast to the compositions shown in Fig. 2, LaSC, SCCr, SCFe, SCAl, SCGa, SCTi, SCZr, SCSn, SCV and SCNb did not contain the hexagonal phase. These samples mainly consist of the perovskite or perovskite related phase. Fig. 3 shows the XRD patterns of these oxides. The XRD patterns of SCFe, SCAl and SCGa show that these samples mainly contain Sr2Co2O5, JCPDS34-1475, orthorhombic phase. In Fig. 3, the peaks corresponding to the orthorhombic phase are marked with “⁎”, and the indices of the diffraction peaks on the basis of the orthorhombic phase are shown above the peaks of SCFe. The orthorhombic phase has the Ca2Fe2O5-type (brownmillerite) structure, which can be explained as oxygen-deficient perovskite with vacancy ordering along the [101]c direction. This oxygen vacancy ordering results in superstructure: ao
bo
co cðM2Þac
4ac
ðM2Þac [26,27].

Fig. 3. X-ray diffraction patterns after sintering at 1200 °C for 5 h in air for LaSC, SCCr, SCFe, SCAl, SCGa, SCTi, SCZr ,SCSn ,SCV and SCNb. Peaks shown by “⁎” correspond to Sr2Co2O5; “sz” to SrZrO3; “ss” to SrISnO3; “v” to Sr3 (VO4)2. Inscriptions above the peaks for LaSC, SCFe and SCNb are the Miler's indexes (hkl) of the cubic perovskite, the orthorhombic brownmillerite and the tetragonal superstructure where at × at × ct ≈ ac × ac × 2ac, respectively.

T. Nagai et al. / Solid State Ionics 177 (2007) 3433–3444

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diffraction patterns, and their positions agree well with those of (020)o, (121)o, (161)o and (222)o peaks of the orthorhombic Sr2Co2O5. These peaks are also shown by “⁎” in Fig. 3. Furthermore, (200)c, (211)c and (220)c peaks of SCNb exhibit line splitting, which cannot be explained by the simple cubic perovskite. The additional peaks and the peak splitting indicate that the oxygen vacancies in the perovskite phase are not disordered completely and the vacancy ordering appears locally [1,3]. In addition to the brownmillerite structure, tetragonal superstructure, whose unit cell is in relation of at
at
ct cac
ac
2ac ; can also appear in the Co-based perovskite upon the oxygen vacancy ordering [27]. The four additional peaks, which seemed to correspond to (020)o, (121)o, (161)o and (222)o of the brownmillerite, originate from the unit cell doubling upon the vacancy ordering: they can also be explained by the tetragonal superstructure as (001)t, (101)t, (103)t and (201)t, respectively. Furthermore, the line splitting in SCNb in Fig. 3 can also be explained by the tetragonal superstructure. Thus the structure of LaSC, SCCr, SCTi, SCZr, SCSn, SCV and SCNb is estimated to be tetragonal with oxygen vacancy ordering rather than cubic perovskite with disordered vacancies. The lattice parameters of SCNb are at = 3.878, ct = 7.799. The indices of the diffraction peaks on the basis of the tetragonal superstructure are shown above the peaks of SCNb in Fig. 3. In the patterns of LaSC, SCCr, SCFe, SCAl, SCGa SCTi and SCNb, there are no additional peaks, which suggest the formation of the second phase. On the other hand, the peaks of the second phase appear in the patterns of SCZr, SCSn and SCV: those are assigned as SrZrO3 (JCPDS 10-268), SrSnO3 (JCPDS22-1442), and Sr3(VO4)2, (JCPDS 29-1318), respectively. The appearance of Sr3(VO4)2 in SCV agrees with the results of Kharton [7]. These results show that the concentration limits of Zr, Sn and V to form the single-phase solid solution with SC are less than 10%. SC has the cubic perovskite structure with a disordered oxygen vacancy arrangement at a higher temperature than 900 °C and it undergoes phase transition to the hexagonal phase with the BaNiO3-type structure when it is cooled below the critical temperature [3]. The above mentioned results suggest that partial substitution with La, Cr, Fe, Al, Ga, Ti, Zr, Sn, Vand Nb stabilizes the high-temperature phase of SC and the phase transition to the low-temperature phase is retarded, but the substitution with Ni, Cu, Zn, In and Ce does not have the same effect. Next, the stability of the perovskite under the condition near the oxygen generator was investigated. The powder samples with perovskite structure were subjected to annealing in oxygen flow at 850 °C for 24 h. Fig. 4 shows the XRD patterns of the samples. All of the diffraction peaks in the XRD patterns of SCCr, SCAl, SCGa, SCZr, SCSn and SCV, can be assigned as SrCoO2.52 hexagonal and additional phases. No peaks corresponding to the perovskite remain in the patterns. These results show that the perovskite phases of these solid solutions are insufficiently stable to remain under the annealing condition, and they completely transform to the BaNiO3-type hex-

Fig. 4. X-ray diffraction patterns after annealing at 850 °C for 24 h in oxygen flow. Peaks shown by “c” correspond to Co3O4 (JCPDS 43-1003); “r” to SrCrO4 (JCPDS 15-356). Ih/Imax, peak intensity ratio of (101) peak for SrCoO2.52, is shown above the peak for LaSC, SCFe, SCTi and SCNb.

agonal phase. In contrast, the XRD patterns of LaSC, SCFe, SCTi and SCNb in Fig. 4 show that these samples still consist mainly of the perovskite. Therefore the stabilities of the perovskite phases of LaSC, SCFe, SCTi and SCNb are higher than those of SCCr, SCAl, SCGa, SCZr, SCSn and SCV. The main phase of SCFe is changed from brownmillerite to perovskite by the annealing. Furthermore the additional peaks and the peak splitting upon the vacancy ordering, which appeared in

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the pattern of LaSC, SCTi and SCNb in Fig. 3, disappear in Fig. 4: the main phase of the three mixed conductors are changed from the tetragonal superstructure to the cubic perovskite structure. These structural changes may be due to a decrease in oxygen vacancy upon heat treatment under higher oxygen partial pressure: the concentration of the oxygen vacancy becomes insufficient to form the vacancy-ordered phase. The traces of the hexagonal phase appeared in the diffraction patterns of LaSC, SCFe, SCTi and SCNb. The peak intensity of the hexagonal phase differs in each system. In order to compare the amount of the hexagonal phase, XRD peak intensity ratio, Ih=Imax ¼ fintensity of the ð101Þ peak for SrCoO2:52 g =ðintensity of the most intense peak in the diffraction patternÞ;

was calculated. The (101) peak for SrCoO2.52 is the second intense peak of the hexagonal, while the most intense peak of the phase is overlapped by the main peak, (110), of the perovskite. The overlapped peak appears as the most intense peak in the diffraction patterns at around 2θ = 32°. According to JCPDS 40-1018, the relative intensity of the (101) peak for SrCoO2.52 is 73%. Thus Ih/ Imax can be changed between 0% and 73% for the perovskiteSrCoO2.52 hexagonal dual phase system. In Fig. 4, the Ih/Imax for the four kinds of solid solutions is shown above the (101) peak of each pattern. The order of Ih/Imax is LaSCN SCFe N SCTi N SCNb. This result shows that the amount of the hexagonal phase is the smallest in SCNb. By changing the annealing temperature from 850 °C, the amount of the hexagonal phase was changed. Fig. 5 shows the dependence of Ih/Imax for LaSC, SCFe, SCTi and SCNb upon the annealing temperature. Annealing was carried out at each temperature for 24 h in oxygen flow. For LaSC after annealing at 700 °C, the perovskite phase completely transformed to the hexagonal phase and Ih/Imax reached 77.7%. At other annealing temperatures, the main phase was the cubic phase with the perovskite structure and Ih/Imax was smaller than 15%. Ih/Imax for LaSC decreases monotonously by increasing the temperature. For SCFe, SCTi and SCNb, the main phase was the perovskite at

Fig. 6. XRD pattern of SCNb after annealing at 900 °C for 200 days in oxygen flow.

any annealing temperature. Ih/Imax reaches maximum at 800 °C but the peak value is less than 6%. At 900 °C, the ratio is decreased to less than 1% for the three solid solutions. The order of Ih/Imax is kept as LaSC N SCFe N SCTi N SCNb at any annealing temperature. The above-mentioned results imply that the perovskite structure of LaSC, SCFe, SCTi and SCNb is not completely stable at 850 °C or below, and the tendency to transform to the hexagonal remains: the tendency of the solid solutions is in the order of Ih/Imax. On the other hand, perovskite seems to be stabilized at 900 °C. In commercial use, the membrane in the oxygen generator will be operated continuously for months. The oxygen permeability of the membrane should not degrade within that period, and the operational conditions must be chosen to avoid the phase transition. In order to examine the stability of the perovskite structure at 900 °C, SCNb was annealed at 900 °C in oxygen atmosphere for 200 days. The XRD pattern of the annealed sample is shown in Fig. 6. After this period of annealing, no traces of the hexagonal phase appear in the pattern. This result indicates that the perovskite of SCNb is sufficiently stable to be applied for the oxygen generator when it is operated at 900 °C. In the figure, small peaks that agree with the tetragonal superstructure exist as shown by “⁎”. The appearance of the vacancy-ordered phase, which was absent after annealing at 850 °C, may be due to an increase in the oxygen vacancy upon the annealing at higher temperature. In the present study, it became clear that the stability of the perovskite structure and the tendency of the transition to the BaNiO3-type structure are changed not only by the annealing temperature but also by the kind of substitutional cation. The stability sequence of the perovskite structure in the SC-based mixed conductor upon the substituting cation is summarized as SCNi; SCCu; SCZn; SCIn; SCCebSCCr; SCAl; SCGa; SCZr; SCSn; SCV bLaSC; SCFe; SCTi; SCNb:

Fig. 5. Dependence of XRD peak intensity ratio, Ih/Imax, of LaSC, SCFe, SCTi and SCNb on the annealing temperature. Annealing was carried out at each temperature for 24 h in oxygen flow.

Furthermore, the stability of the four kinds of the solid solutions is ordered as LaSC b SCFe b SCTi b SCNb by referring to the tendency of the transition to the hexagonal phase. These findings indicate that SCNb has the highest stability of the perovskite phase among the samples examined in this study.

T. Nagai et al. / Solid State Ionics 177 (2007) 3433–3444 Table 1 Tolerance factor (t) and ionic radii (r) of SC-based mixed conductors t

r

SCCu SC

0.993 0.999

SCAl SCCr SCGa LaSC SCFe SCTi SCNb

1.003 0.999 0.999 0.995 0.997 0.999 0.998

Cu2+: 0.73 Sr2+: 1.44 Co3+: 0.61 Al3+: 0.53 Cr3+: 0.615 Ga3+: 0.62 La3+: 1.32 Fe3+: 0.645 Ti4+: 0.605 Nb5+: 0.64

Tolerance factor was calculated using ionic radii of ions by Shannon et al. [31]: the valence of the B-site cations were assumed as Cu2+, Al3+, Cr3+, Ga3+, La3+, Fe3+ with high spin, Co3+ with high spin, Ti4+ and Nb5+, respectively. The effect of the oxygen vacancy was ignored in the calculation.

The above results show that the effect of the stabilization of the perovskite structure differs even in cations having the same valence, i.e. the order of the effect is In3þ bfCr3þ ; Al3þ ; Ga3þ g; Ce4þ bfZr4þ ; Sn4þ gbTi4þ ; V5þ bNb5þ : It should be noted that In3+, Ce4+, Zr4+, Sn4+ and V5+, which have a lower stabilization effect, have a solubility lower than 10% for SC. Thus, it is suggested that cations with lower solubility to SC exhibit a lower effect on the stabilization of the perovskite structure. Consequently, cations that could not form the mono-phase solid solution should be eliminated from the discussion on the relation between the valence of the substitutional cation and the stability of the perovskite. In this case, the sequence of the stabilization effect is Cu2þ bCr3þ ; Al3þ ; Ga3þ bLa3þ ðsubstituting A  siteÞb 1

Fe3:108þ bTi4þ bNb5þ : This shows that the stability of the perovskite structure is enhanced by substituting the B-site with a cation having a higher valence, as predicted. Tolerance factor is known to largely control the stability of the perovskite [28–30]. Here the relation between stability of the perovskite structure in the SC-based mixed conductors and tolerance factor is discussed. The factor is defined as t ¼ ðrA þ rO Þ=M2ðrB þ rO Þ; where rA, rB and rO correspond to the average ionic radii in Asite, the one in B-site and ionic radii of the oxygen ion, respectively. The tolerance factor was calculated using ionic radii by Shannon et al. and the results are shown in Table 1 [31]. A set of A- and B-site cations where t = 1 is thought to correspond to an ideal perovskite. Where t N 1, the A-site cations are too large to form a perovskite structure and a hexagonal structure like BaNiO3-type tends to appear. As shown in the table, an ionic radius of Sr2+ is larger than La3+. Therefore, Sr substitution for LaCoO3 increases the tolerance factor and it 1

By thermogravimetric measurement, the average valence of Co and Fe in SCFe was estimated to be 3.108 at 850 °C in oxygen. See Table 2.

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reaches almost unity at SrCoO3–δ. This change qualitatively explains the appearance of BaNiO3-type hexagonal in SC and the perovskite stabilization effect of La substitution for SC. The sequence of the tolerance factor is SCCubLaSCbSCFebSCNbbSC; SCCr; SCGa; SCTibSCAl: This result suggests that the stabilization of the perovskite structure upon Fe and Nb substitution can be explained by reduction of the tolerance factor. But the sequence of the factor for LaSC, SCFe and SCNb is different from that of the effect of the stabilization for the perovskite structure. The factor of SCCu is the smallest in the table and this implies the stabilization effect of Cu substitution for SC; however, Cu substitution had no effect on stabilization. Furthermore, stabilization of the perovskite upon substitution by Cr, Ga, Ti and Al cannot be explained by the factor. It is difficult to understand the stability sequence of the perovskite structure only by the sequence of the tolerance factor. In this study, the stability of the perovskite structure in the SCbased mixed conductors was found to have a close relation to the valence of the substituting cation. This implies that change in the oxygen nonstoichiometry, 3–δ, upon cation substitution should be closely connected to the perovskite stability. The effect of the cation substitution on the oxygen nonstoichiometry can be qualitatively discussed using a simple point defect model expressed in Kröger– Vink notation. LaCoO3 is chosen as a standard reference state and assumed to be defect free. Sr substitution for LaCoO3 will be compensated through either one or both of the following ways:
V SrO þ La
La þ CoCo þ ð1=4ÞO2 ðgÞ Z SrLa þ CoCo þ LaO3=2 ;

d

dd


V 2SrO þ 2La
La þ OO Z 2Sr La þ VO þ 2LaO3=2 ;

··

where V O represents doubly ionized oxygen vacancy and CoCo = Co4+. Thus the defect reaction for SC can be written as:

·

V V
þ CoCo þ 3O
SrLa O () Sr La þ ð1−2dÞCoCo þ 2dCoCo

d

dd

þ ð3−dÞO
O þ dVO þ ðd=2ÞO2 ðgÞ;

d

where δ shows the concentration of the oxygen vacancy. This result shows that both oxygen vacancy and Co4+ may exist in SC: 3–δ and the amount of Co4+ will change with oxygen partial pressure. The substitution of SC by cation with higher valence than Sr2+ and Co3+, which is donor doping for SC, will result in another defect compensation. In the case of Nb substitution, the defect compensation will be either or both of the following:

dd

NbO5=2 þ 2CoCo Z NbCo þ Co
Co þ CoO3=2 þ ð1=2ÞO2 ðgÞ;

d

dd

dd


NbO5=2 þ Co
Co þ V Z NbCo þ OO þ CoO3=2 :

The defect reaction for SCNb can be written as: V
V SrLa þ 0:8CoCo þ 0:1Co
Co þ 0:1 NbCo þ 3OO () SrLa

dd

d

dd


þ ð0:8−2dÞCoCo þ ð0:1þ2dÞCo
Co þ 0:1 NbCo þ ð3−dÞOO

dd

d

þ dVO þ ðd=2ÞO2 ðgÞ:

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T. Nagai et al. / Solid State Ionics 177 (2007) 3433–3444

Table 2 3–δ at initial condition (3–δ0), 3–δ at 850 °C in oxygen (3–δ850,O2) and average valence of cation in B-site at 850 °C in oxygen (Z850,O2)

SCCu SC SCAl LaSC SCFe SCTi SCNb

3–δ0

3–δ850,O2

Z850,O2

2.637 2.737 2.537 2.804 2.723 2.735 2.773

2.493 2.507 2.539 2.614 2.554 2.568 2.617

2.986 3.014 3.078 3.128 3.108 3.136 3.234

3–δ0 was measured at 80 °C after heat treatment at 1000 °C for 0.5 h in air flow. 3–δ850,O2 was measured during the cooling step from annealing at 1000 °C for 0.5 h in oxygen flow. Z850,O2 was estimated using 3–δ850,O2.

Therefore donor doping for SC will result in a decrease of oxygen vacancy and/or Co4+. Change in oxygen vacancy and Co4+ can be distinguished by measuring oxygen nonstoichiometry, 3–δ. The average valence of B-site cation can be determined from 4–2δ for SC and B-site substituted mixed conductors. 3–δ for the SC-based mixed conductors were measured by thermogravimetry. Table 2 shows 3–δ of the mixed conductors at initial condition, 3–δ0; the one at 850 °C in oxygen, 3–δ850, O2; and average valence of cation in B-site at 850 °C in oxygen, Z850,O2. Since SC and SCCu changed from the perovskite structure to the SrCoO2.52 hexagonal structure between 800 °C and 750 °C under the TG-DTA condition, 3–δ850,O2 for SC and SCCu correspond to their cubic phase. 3–δ850,O2 for SC was reported by Takeda et al. [22]: from the line of x for procedure I in O2 gas in Fig. 1 of the paper, their result of 3–δ850,O2 for SC is estimated as 2.5. This result agrees well with ours, 2.507. The sequence of 3–δ850,O2 is

electrostatic repulsion and destabilize the hexagonal structure: consequently the substitution acts as stabilizing the perovskite structure. As shown in Table 2, Z850,O2 for LaSC is larger than SCFe but the stability of the perovskite structure in LaSC was lower than SCFe: LaSC seems to be an exception for the above discussion. This will be explained as follows; the A-site cations are sited out of the chains of the face-shared oxygen octahedral in the hexagonal; an increase in the valence of A-site cation may change the valence of B-site cation but the repulsion between the substituted La3+ and other cations does not largely contribute to increase the electrostatic loss. Therefore the efficiency of the A-site substitution for the stabilization of perovskite will be lower than that of the B-site substitution. In this study, SCNb, SCTi, SCFe and LaSC exhibited higher stability of the perovskite phase. Next, ceramic samples of the solid solutions, as well as SC and Sr(Co0.8Fe0.2)O3–δ (SCF1082) were prepared and the oxygen permeability of the samples was measured. Fig. 7(a) shows the change in the oxygen permeability of SC with temperature. Oxygen permeability of SC exhibited large hysteresis with temperature. In the course of heating, the oxygen permeability of SC is extremely low below 850 °C, but it jumped by about two orders of magnitude at 900 °C. In cooling between 950 °C and 750 °C high oxygen

SCCubSCbSCAlbSCFebSCTibLaSCbSCNb; while the one for Z850,O2 is SCCubSCbSCAlbSCFebLaSCbSCTibSCNb: These results suggest that, except for LaSC, a rise in the valence of the substitutional cation increases both 3–δ850,O2 and Z850,O2, and enhances the stability of the perovskite structure. Here we would like to discuss a feature of the BaNiO3-type hexagonal in order to understand the mechanism of the stabilization of perovskite structure upon the cation substitution. Goodenough et al. and Takeda et al. pointed out that the distance between B-site cations, which have the highest valence in SC-based oxides, is shorter than perovskite structure because the hexagonal structure has face-shared oxygen octahedral while the perovskite structure has corner-shared one. Owing to the crystal-chemical character, large electrostatic repulsion between the B-site cations can make the hexagonal more unstable than the perovskite structure. In spite of this, the hexagonal structure can be competitive upon the counteraction of the repulsive energy by the increase of covalency in the B–O bond and by the shielding effect of oxygen on the B-site cations [29,32]. The above discussion leads to a prediction that the substitution with higher valence cation and increasing the average valence in B-site of SC-based oxides raises the

Fig. 7. Dependence of oxygen permeability on temperature: SC (a); SCNb (b).

T. Nagai et al. / Solid State Ionics 177 (2007) 3433–3444 Table 3 Sintering temperature of the ceramic samples (T), density (ρ) and relative density of the ceramic samples (ρre), gas leak of the samples detected during the permeability measurement (Leak) and oxygen permeability (Jo2) at 850 °C and 900 °C of SC, LaSC, SCFe, SCTi, SCNb and SCF1082

SC LaSC SCFe SCTi SCNb SCF1082

T (°C)

ρ (g/cm3)

ρre (%)

Leak

Jo2 (cm3/min/cm2) at 850 °C

at 900 °C

1200 1200 1200 1250 1250 1250

5.15 5.27 5.10 5.08 5.28 5.17

87.8 92.1 96.0 94.9 96.0 95.6

Small Medium Large Not detected Not detected Small

3.46 2.77 3.21 3.05 3.50 1.30

4.21 3.36 3.93 4.10 4.24 1.62

The oxygen permeability was measured during the cooling step: the sample thickness is 1 mm: oxygen activity of the sample surface is 0.21 atm and 0.01 atm. The permeability is shown in a volumetric unit under standard temperature pressure condition. Leak “small” means the helium leak in the outlet gas was less than 100 ppm; “medium” means the one was between 100 ppm and 500 ppm; “large” means the one was larger than 500 ppm.

permeability was maintained. Since the low oxygen permeable hexagonal phase is believed to be stable below 850 °C in SC, the high-oxygen permeability observed during cooling below 850 °C should not be stable. In spite of this prediction, the oxygen permeability was stable during the measurement for dozens of minutes even below 850 °C. Fig. 7(b) shows the change in the oxygen permeability of SCNb with temperature. The oxygen permeability shows monotonous increase and decrease with temperature. There is no large hysteresis in permeability, but the permeability during heating and cooling is not the same. The permeability of LaSC, SCFe, SCTi and SCF1082 also showed monotonous temperature dependence like SCNb. For SC, LaSC, SCFe, SCTi, SCNb and SCF1082, sintering temperature, density, relative density, gas leak property detected during the permeability measurement and the oxygen permeability at 850 °C and 900 °C are shown in Table 3. The theoretical density was estimated by using the lattice constant and the value of 3–δ Except for SCF1082, 3–δ0 in Table 2 was adopted as 3–δ. For SCF1082, 3–δ at room temperature in air was estimated from the 3–δ–T diagram, shown in Fig. 1 of Ref. [33]. Samples densified greater than 90% to the theoretical density were successfully prepared for the solid solutions, but the maximum density of the ceramic SC sample was 87.8%. In addition, cracks were visible to the naked eye on the surface of the ceramic samples of SCFe. Similar cracks were also visible on LaSC, but the number of cracks was greater for SCFe than LaSC. During the measurement, oxygen permeability reached equilibrium faster in the cooling step than the heating step. Thus permeability measured during the cooling step is estimated to be more reliable. Oxygen permeability shown in Table 3 is the value measured during the cooling step. The data are normalized to the following conditions; the sample thickness is 1 mm, oxygen partial pressure of the sample surface is 0.21 atm and 0.01 atm. During the oxygen permeability measurement, contaminating helium was detected in the sweeping argon gas for SC, LaSC, SCFe and SCF1082. This indicates that the prepared disc samples

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of these compositions have a gas leak pass. In contrast, no gas leak was detected for SCTi and SCNb. Consequently, no correction due to gas leak was necessary for evaluating the oxygen permeability of SCTi and SCNb, while correction was necessary for SC, LaSC, SCFe and SCF1082. It is strange that the disc of SCFe showed the largest gas leak although the relative density of the sample was the highest of all the samples. One of the reasons for the large gas leak of SCFe may be the number of cracks in the disc sample as mentioned above. In this study, the permeability of SC, LaSC, SCFe, SCTi and SCNb is more than double that of SCF1082 as shown in Table 3. Among the samples, SCNb and SC showed the highest permeability at 850 °C and 900 °C. The order of permeation is SCNbzSCNSCFeNSCTiNLaSCNSCF1082; at 850-C; SCNbzSCNSCTiNSCFeNLaSCNSCF1082; at 900-C: In contrast, the stability of the perovskite structure was found to be ordered as SCNbNSCTiNSCFeNLaSCNSC: As discussed above, the stability of the perovskite structure in the SC-based solid solutions should be in relation to the valence of the substitutional cation and oxygen nonstoichiometry. This consideration may lead to the conclusion that the oxygen permeability, which is controlled by the concentration of the oxygen vacancy, is in a trade-off-relation against the stability of the perovskite structure. However, the above order of oxygen permeability does not seem to correlate with the order of the stability of the perovskite structure: this suggests that the oxygen permeability of the mixed conductor is not necessarily in a trade-off-relation against the stability of the perovskite phase. There are discrepancies in the oxygen permeability data reported by Teraoka et al. [2,3] and those reported by other groups. Such conflicting results reflect the difficulties in measuring the oxygen fluxes at high temperatures. Bouwmeester and Burggraaf pointed out that the oxygen permeation through perovskite mixed conductors is changed by measurement conditions, which may be the reason for the discrepancies [1]. For reference, the permeability data of SCF1082 from the literature are listed and compared with the present results in Table 4 [2,5,16,34]. Our findings are less than half of the findings of Teraoka et al., and are similar to the findings of Huang et al. The discrepancy in the permeability of SCF1082 is still large. However, we believe that comparative discussion of the permeability between the solid solution samples measured under the same conditions is still relevant. Oxygen permeation is sensitive to the conditions of the sample surface [16,35,36]. In the present study, no sample surface modification that is effective in improving the oxygen exchange kinetics at the surface was carried out. Thus, the oxygen permeability of the solid solutions can be increased by treating the sample surface. Further investigation is necessary to clarify the relation between the oxygen nonstoichiometry, the oxygen permeability and the stability of the perovskite structure in mixed conductors.

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T. Nagai et al. / Solid State Ionics 177 (2007) 3433–3444

Table 4 Oxygen permeability of SCF1082 in present result and references Jo2 (cm3/min/cm2)

Conditions

Table 5 Thermal expansion coefficient of LaSC, SCTi, SCNb and SCFe

at at 850 °C 900 °C Present result

1.30

1.62

Teraoka 3.1 et al. Kruidhof 0.23 et al. Huang 1.36 et al.

–

Qiu et al. 0.89

1.06

0.28 –

low L = 1 mm: Pohigh 2 /Po2 = 0.21/0.01 atm: Jo2 low depended on log(Pohigh 2 /Po2 ); Jo2 at the above driving-force was estimated by using the dependence. L = 1 mm: Air/helium flow (30 cm3/min): data from Fig. 3. L = 1 mm: Air/helium flow: data was red from Fig. 3. L = 1.48 mm: Air/helium flow: data was red from Fig. 6 where {(Po2)n–(Po2)n} was (0.210.75–0.010.75) = 0.279; shown data was normalize to L = 1 mm by multiplying 1.48. L = 1 mm: Air/helium flow (40 cm3/min): data was estimated from Fig. 3.

αl (ppm/°C)

Ref.

–

[2]

LaSC SCTi SCNb SCFe

25.9 22.2 22.8 16.7 (in air) 28.4 (in O2)

After keeping at 1000 °C for 30 min, the sample length was measured in the cooling step: the coefficient is for the temperature between 850 °C and 50 °C: the sample atmosphere was airflow except for the oxygen flow measurement of SCFe.

[5] [16]

temperature was measured in the cooling step after keeping at 1000 °C for 30 min. Fig. 8(a) and (b) show the change in the sample length of SCNb and SCFe with temperature, respectively. The y-axis represents relative sample length,

[34]

Conditions: sample thickness (L): oxygen partial pressure at the sample surface low (Pohigh 2 /Po2 ) or atmosphere of each side of the sample: other comments.

Finally, the thermal expansion curve of the solid solutions LaSC, SCFe, SCTi and SCNb was measured. The thermal expansion of the solid solutions was affected by the history of the sample. Therefore, the change in the sample length with

DLðT Þ=LðT0 Þ ¼ fLðT Þ−LðT0 Þg=LðT0 Þ; where L(T) is the sample length at T °C and T0 is 1000 °C. The change in the length of SCNb measured in air, as shown in Fig. 8(a), is not linear but monotonous with temperature. LaSC and SCTi also showed a similar change. In contrast, the length of SCFe measured in air, as shown in Fig. 8(b), exhibits a peak between 600 °C and 400 °C. This anomaly in the contraction curve was reproducible. Fig. 8(b) also shows the contraction curve of SCFe measured in oxygen. In this case, the peak does not appear and the curve is monotonous. These results suggest that the anomaly in the contraction curve of SCFe in air is caused by the increase in the oxygen vacancy. Since SCFe has the brownmillerite structure with vacancy ordering, the anomaly in the curve may be due to the order–disorder transition of the oxygen vacancy [1]. The anomaly in the curve of SCFe may be the origin of the quantities of the cracks in the ceramic sample and large gas leak during the oxygen permeability measurement of the disk sample. The thermal expansion coefficients, αl, of the solid solutions between 850 °C and 50 °C are shown in Table 5. The coefficients of LaSC, SCTi and SCNb are 22 ppm/°C or larger. On the other hand, that of SCFe is as small as 16.7 ppm/°C. This small coefficient is caused by the small change in the sample length below 400 °C, in the lower temperature region than the peak in the contraction curve, as shown in Fig. 8(b). In contrast to this, SCFe showed an extremely high thermal expansion coefficient up to 28.4 ppm/°C in oxygen flow, where the order–disorder transition did not take place. In the present study, SCNb showed the highest stability of the perovskite structure, easiness in preparing gastight dense ceramics, the highest oxygen permeability and monotonous thermal expansion. Nb was found to be effective as a substitutional cation for preparing an SC-based mixed conductor. 4. Conclusions

Fig. 8. Change in the sample length with temperature: SCNb (a); SCFe (b). The measurement was carried out during cooling after keeping at 1000 °C for 30 min in each atmosphere. The y-axis represents the relative sample length change. T0 is 1000 °C.

In this study, the dose of the substitutional cation for SrCoO3–δ, SC, was fixed at 10 mol% of the A- or B-site, and 15 kinds of SC-based oxides, (La0.1Sr0.9)CoO3–δ (abbreviated as LaSC) and Sr(Co0.9X0.1)O3–δ, where X was Ni, Cu, Zn, Cr,

T. Nagai et al. / Solid State Ionics 177 (2007) 3433–3444

Fe, Al, Ga, In, Ce, Ti, Zr, Sn, V and Nb (abbreviated as SC“X”) were synthesized. The stability of the perovskite structure, oxygen nonstoichiometry, oxygen permeation properties and thermal expansion were examined and their dependence upon the substitutional cation was discussed. The results are summarized as follows. 1. Solubility of Ni, Zn, In, Ce, Zr, Sn and V for B-site of SC was less than 10 mol%. A mono-phase solid solution could not be prepared for these compositions. 2. In the case that the stability of the perovskite structure is insufficient, it transforms to the hexagonal SrCoO2.52 with the BaNiO3-type structure. The tendency of the transformation of the perovskite structure was evaluated by the appearance and the amount of the hexagonal phase. By the XRD investigation of the SC-based oxides after sintering and post-annealing in oxygen, the stability sequence of the perovskite structure upon the substitutional cation was expressed as SCNi; SCCu; SCZn; SCIn; SCCebSCCr; SCAl; SCGa; SCZr; SCSn; SCVbLaSCbSCFebSCTibSCNb: 3. Cation with lower solubility to SC exhibited a lower effect on the stabilization of the perovskite structure. 4. As far as a single-phase solid solution was formed, the stability of the perovskite structure was enhanced by substituting in the Co-site with a cation having higher valence. 5. Through thermogravimetric measurement, the oxygen nonstoichiometry, 3–δ, was determined. The sequence of 3–δ at 850 °C in oxygen was SCCubSCbSCAlbSCFebSCTibLaSCbSCNb: For Co-site substitution, 3–δ was increased with valence of the substitutional cation. 6. Stabilization of the perovskite structure upon the cation substitution was correlated to an increase in the electrostatic repulsion between B-site cations and a larger impact of the repulsion on the stability of BaNiO3-type hexagonal structure than perovskite. 7. During the oxygen permeability measurement, the ceramic samples of SCTi and SCNb were gastight. On the other hand, those of SC, LaSC, SCFe and Sr(Co0.8Fe0.2)O3–δ (SCF1082) exhibited a gas leak. The SCFe sample had many cracks and it showed the largest leak. 8. The oxygen permeability of SC exhibited large hysteresis with temperature, while those of SCNb, SCTi, SCFe, SCF1082 and LaSC showed monotonous change with temperature. The order of the oxygen permeability of the ceramic samples at 900 °C was SCNb≥SCNSCTiNSCFeNLaSCNSCF1082: 9. The oxygen permeability of the mixed conductors was not in a trade-off relation against the stability of the perovskite phase.

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10. LaSC, SCTi and SCNb contracted monotonously and their thermal expansion coefficients, αl, between 1000 °C and 50 °C were 23 ppm/°C or larger. SCFe showed a small αl, as 18.7 ppm/°C and the length change of SCFe exhibited a peak between 600 °C and 400 °C during cooling in air. The peak in the contraction curve did not appear in oxygen gas flow. The appearance of the peak was correlated to the order–disorder transition of the oxygen vacancy, which seemed to be the cause of the cracks in the ceramic sample. Acknowledgements The authors wish to acknowledge the financial support of the Ministry of Economy, Trade and Industry, Japan and The Japan Research and Development Center for Metals (JRCM). References [1] H.J.M. Bouwmeester, A.J. Burggraaf, in: P.J. Gellings, A.J. Bouwmeester (Eds.), The CRC Handbook of Solid State Electrochemistry, CRC Press, Boca Raton, 1997, p. 481. [2] Y. Teraoka, H.M. Zhang, S. Furukawa, N. Yamazoe, Chemistry Letters (1985) 1743. [3] Y. Teraoka, T. Nobunaga, N. Yamazoe, Chemistry Letters (1988) 503. [4] Y. Teraoka, H. Shimokawa, H. Kusaba, K. Sasaki, in: Philippe Knauth, Christian Masquelier, Enrico Traversa, Eric D. Wachsman (Eds.), SolidState Ionics—2004, Mater. Res. Soc. Symp. Proc., vol. 835, Materials Research Society, Warrendale, PA, 2005, p. 77. [5] H. Kruidhof, H.J.M. Bouwmeester, R.H.E.V. Doorn, A.J. Burggraaf, Solid State Ionics 63–65 (1993) 816. [6] V.V. Kharton, L. Shuangbao, A.V. Kovalevsky, E.N. Naumovich, Solid State Ionics 96 (1997) 141. [7] V.V. Kharton, A.V. Kovalevsky, A.P. Viskup, E.N. Naumovich, O.P. Reut, Proceedings of the Third International Symposium on Ionic and Mixed Conducting Ceramics, 31 Aug.–5 Sept., Paris, France, Electrochemical Society Proceedings, vol. 97–24, 1997, p. 736. [8] A.L. Shaula, V.V. Kharton, N.P. Vyshatko, E.V. Tsipis, M.V. Patrakeev, F.M.B. Marques, J.R. Frade, Journal of the European Ceramic Society 25 (2005) 489. [9] Z. Shao, W. Yang, Y. Cong, H. Dong, J. Tong, G. Xiong, Journal of Membrane Science 172 (2000) 177. [10] Z.Q. Deng, W. Liu, C.S. Chen, H. Lu, W.S. Yang, Solid State Ionics 170 (2004) 187. [11] C.G. Fan, Z.Q. Deng, Y.B. Zuo, W. Liu, C.S. Chen, Solid State Ionics 166 (2004) 339. [12] V.V. Vashook, M. Al Daroukh, H. Ullmann, Ionics 7 (2001) 59. [13] N.E. Trofimenko, J. Paulsen, H. Ullmann, R. Müller, Solid State Ionics 100 (1997) 183. [14] J. Tong, W. Yang, B. Zhu, R. Cai, Journal of Membrane Science 203 (2002) 175. [15] L. Yang, X. Gu, L. Tan, L. Zhang, C. Wang, N. Xu, Separation and Purification Technology 32 (2003) 301. [16] K. Huang, J.B. Goodenough, Journal of Electrochemical Society 148 (2001) E203. [17] J.W. Stevenson, T.R. Armstrong, R.D. Carneim, L.R. Pederson, W.J. Weber, Journal of Electrochemical Society 143 (1996) 2722. [18] H.U. Anderson, Solid State Ionics 52 (1992) 33. [19] P.N. Dyer, R.E. Richards, S.L. Russek, D.M. Taylor, Solid State Ionics 134 (2000) 21. [20] J.C. Grenier, L. Fournès, M. Pouchard, P. Hagenmuller, Material Research Bulletin 21 (1986) 441. [21] United States Patent No. US 6875528. [22] Y. Takeda, R. Kanno, T. Takada, O. Yamamoto, M. Takano, Y. Bando, Zeitschrift fuer Anorganische und Allgemeine Chemie 540/541 (1986) 259. [23] W.T. Harrison, S.L. Hegwood, A.J. Jacobson, Journal of the Chemical Society Chemical Communications (1995) 1953.

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