A CO2‐Tolerant Perovskite Oxide with High Oxide Ion and Electronic Conductivity

Mixed ionic–electronic conductors (MIECs) that display high oxide ion conductivity (σo) and electronic conductivity (σe) constitute an important family of electrocatalysts for a variety of applications including fuel cells and oxygen separation membranes. Often MIECs exhibit sufficient σe but inadequate σo. It has been a long‐standing challenge to develop MIECs with both high σo and stability under device operation conditions. For example, the well‐known perovskite oxide Ba0.5Sr0.5Co0.8Fe0.2O3−δ (BSCF) exhibits exceptional σo and electrocatalytic activity. The reactivity of BSCF with CO2, however, limits its use in practical applications. Here, the perovskite oxide Bi0.15Sr0.85Co0.8Fe0.2O3−δ (BiSCF) is shown to exhibit not only exceptional bulk transport properties, with a σo among the highest for known MIECs, but also high CO2 tolerance. When used as an oxygen separation membrane, BiSCF displays high oxygen permeability comparable to that of BSCF and much higher stability under CO2. The combination of high oxide transport properties and CO2 tolerance in a single‐phase MIEC gives BiSCF a significant advantage over existing MIECs for practical applications.

Mixed ionic-electronic conductors (MIECs) that combine the required intrinsic transport properties of high oxide ion conductivity (σ o ) and electronic conductivity (σ e ) with high stability under device operation conditions are needed to accelerate the commercialization of a variety of electrochemical technologies. For example, oxy-fuel combustion coupled with MIEC oxygen separation membranes offers an economically viable solution for CO 2 capture and storage (CCS) for fossil-fuel fired power plants. [1,2] This technology involves combustion of fuels in pure The σ o of BSCF calculated from the oxygen diffusion coefficient D* is 0.07-0.17 S cm −1 at 600 °C [7][8][9] and 0.24 S cm −1 at 700 °C, [8] among the highest for known MIEC perovskite oxides, [7][8][9] giving rise to remarkably high oxygen permeability when used in oxygen separation membranes. [10] However, MIECs with solely large alkaline earth ions such as Ba and Sr on the A-site (e.g., BSCF) have high basicity and have a tendency to react with CO 2 , leading to a degradation in permeation performance. [13,14] Despite the exceptional σ o and electrocatalytic performance of BSCF, this instability toward CO 2 limits application in oxygen separation membranes, especially in environments where CO 2 is prevalent such as in oxyfuel combustion. Various dopants have been employed to increase CO 2 tolerance of BSCF but typically at the cost of reduced oxygen permeability. [2] Here, the single-phase perovskite oxide Bi 0.15 Sr 0.85 Co 0.8 Fe 0.2 O 3-δ (BiSCF) is shown to retain the exceptional oxide transport properties of BSCF and simultaneously exhibit a high CO 2 tolerance. The Bi 1−x Sr x Co 1−y Fe y O 3−δ phases were selected based on a number of hypotheses below. First, the 6s 2 electronic configuration gives Bi 3+ a large Shannon dielectric polarizability (6.12 Å 3 ) [12] and a flexible oxide coordination chemistry supporting many BiO geometries [15] with low BiO bond strengths, [16] which have been shown to enable high oxygen mobility in a number of perovskite oxides. [17][18][19][20][21][22][23][24] Second, Bi doping is expected to increase CO 2 tolerance due to reduced basicity of Bi 3+ oxides when compared against BSCF. Thermogravimetric (TGA) measurements under a CO 2 atmosphere show that BaCO 3 decomposes at ≈1340-1360 °C [25] whereas the decomposition of Bi 2 O 2 CO 3 begins at ≈400-430 °C, reflecting the reduced thermal stability of Bi 3+ versus alkaline earth carbonates. [26] Third, bismuth surface enrichment has been observed in some bismuth-containing oxides. [27][28][29] It is hypothesized that a surface enriched in the less basic Bi 3+ will also reduce reactivity to CO 2 . These potential advantages of high oxygen mobility and reduced basicity lead to the exploration of Bi 1−x Sr x Co 1−y Fe y O 3−δ phases with the aim of coupling high σ o with high CO 2 tolerance in a perovskite MIEC.
There are to date only a few reports in the literature on the Bi 1−x Sr x Co 1−y Fe y O 3−δ family. [30][31][32][33][34] The series of Bi 0.15 Sr 0.85 Co 1−x Fe x O 3−δ (0 ≤ x ≤ 1.0) perovskite oxides were previously synthesized by quenching in air, with measurement of structural and magnetic properties, but no transport properties were reported. [35,36] In contrast to the quenching method used in the previous report, [36] all Bi 0.15 Sr 0.85 Co 1−x Fe x O 3−δ (0.2 ≤ x ≤ 1.0) samples in the present work were furnace-cooled in air with a cooling rate of 5 °C min −1 . The X-ray diffraction (XRD, Figure S3, Supporting Information) and neutron powder diffraction (NPD, Figures S4 Figure 1a. The lattice parameters (a) at room temperature lie in the range 3.8838-3.9013 Å ( Figure S3b, Supporting Information), close to the values reported previously [36] and smaller than that of BSCF (3.9730 Å, PDF Card No. 00-055-0563). As such, the AO distances are shorter in BiSCF (2.7231 Å) than in BSCF (2.8228 Å). This will stabilize the oxide over the carbonate, consistent with the observation that larger cations stabilize salts with larger anions (carbonate Adv. Mater. 2020, 32,1905200  in this case) on ionic model lattice enthalpy grounds. [37] The bond valence sum for Sr 2+ in BiSCF is 2.282 and is 2.116 for Bi 3+ which are not too dissimilar from the expected 2.15 valence for the A site. This is contrasted with a larger average A site bond valence sum of 2.315 in BSCF and significant disparity between Sr 2+ (1.786) and Ba 2+ (2.843). The incorporation of Bi 3+ in BiSCF introduces covalent bonding between A-site cations and oxygen, which reduces the ionic character of the AO bond when compared with BSCF.
BiSCF exhibits considerable oxygen nonstoichiometry, with δ = 0.39(1) at room temperature and ambient oxygen partial pressure, which is significantly larger than that in LSCF (δ < 0.05) [3,4] and slightly lower than that of BSCF (δ ≈ 0.45) [38] measured under similar conditions. The oxygen content (3−δ) of BiSCF at high temperatures was determined using neutron diffraction, cerimetric titration, and TGA which all give consistent data (Table S4, Supporting Information). Upon heating from 500 to 800 °C, BiSCF releases oxygen and the oxygen content decreases to 2.466(9) at 800 °C ( Figure 1b and Table S4, Supporting Information). This is a result of the temperature-and oxygen partial pressure (pO 2 )-dependent equilibrium of the oxygen sublattice with the gas phase, where higher temperature entropically favors dioxygen gas, and is a common characteristic of perovskite oxides containing transition metals with variable valences such as Fe and Co. [4,39] The Bi 0.15 Sr 0.85 Co 1−x Fe x O 3−δ materials exhibit high total electrical conductivity (σ), which varies with temperature and Co/Fe ratio ( Figure 1c). The compositions with higher Co contents exhibit higher σ. In particular, the Bi 0.15 Sr 0.85 Co 0.8 Fe 0.2 O 3−δ (BiSCF) composition shows a maximum σ of ≈365 S cm −1 at ≈500 °C, and σ remains close to 200 S cm −1 up to 800 °C. The maximum σ of BiSCF is comparable to that of LSCF (≈350-400 S cm −1 ) [3,4] and significantly higher than that of BSCF (≈50 S cm −1 , Figure 1c).
BiSCF also exhibits exceptionally high oxygen diffusivity, as revealed by 18 O tracer measurements using ceramic samples and the secondary ion mass spectrometry (SIMS)-line scan technique (Figure 2). The oxygen tracer diffusion coefficient, D*, provides a direct and quantitative description of the oxygen transport properties. The D* values of BiSCF are ≈1.4 × 10 −7 and 7.2 × 10 −7 cm 2 s −1 at 600 and 700 °C, respectively ( Figure 2a). These D* values are among the highest for the known perovskite oxides and compare well with those of BSCF and other MIECs (Figures S10 and S11, Supporting Information). For example, the D* of BiSCF is approximately three orders of magnitude higher than that of LSCF at 600 °C (D* is ≈2.2 × 10 −10 at 594 °C for LSCF). [40] For the data collected using the line scan technique, the D* value of BiSCF is slightly lower than that of BSCF at temperatures ≤600 °C (D* values of BSCF ceramic [7] and thin film [8] samples are 3.0 × 10 −7 to 3.6 × 10 −7 cm 2 s −1 at ≈600 °C), but increases more rapidly with temperature and becomes higher than that of BSCF above 700 °C. The activation energy, E a , for tracer diffusion in BiSCF is 1.2 eV, which is higher than the 0.5-1.1 eV for BSCF, [7,8] but lower than the ≈1.9 eV of LSCF. [40] [7] and La 0.6 Sr 0.4 Co 0.2 Fe 0.8 O 3−δ (LSCF) [40] measured with ceramic samples using the SIMS-line scan technique. The exchange oxygen partial pressure (pO 2 ) is ≈0.21 bar for all ceramic samples including those from the literature. Data for BSCF ceramic samples measured using SIMS-depth profile within individual single grains [7] and BSCF thin film samples (exchange pO 2 is 0.5 bar) measured using SIMS-line scan [8] are also presented. Comparison with other known MIECs including La 0.3 Sr 0.7 CoO 3−δ (LSC), GdBaCo 2 O 5+δ (GBCO) and La 2 NiO 4+δ (LNO) is shown in Figures S10 and S11 (Supporting Information). Typical diffusion profiles and fits for isotope diffusion profiles at 600 and 650 °C are shown in Figure S12 (Supporting Information).
Since the oxide ion conduction is achieved via the migration of oxygen vacancies, the oxygen vacancy diffusion coefficient, D v , is calculated from the measured D* and oxygen content using Equations (S3) and (S4) (Supporting Information), and compared with other MIECs (Figure 2c and Figure S11a, Supporting Information). D v of BiSCF is higher than that of LSCF (by 1-2 orders of magnitude at 550-750 °C) and compares favorably with that of BSCF above 600 °C (D v at 700 °C = 5.9 × 10 −6 cm 2 s −1 for BiSCF and 3.2 × 10 −6 cm 2 s −1 for BSCF [8] ).
The σ o value, calculated via the Nernst-Einstein relationship (Equation (S5), Supporting Information) using the measured D*, a correlation factor of 0.69 (for a cubic perovskite), [41] the oxygen content at the measurement temperature (Table S4, Supporting Information) and cell volume ( Figure S3b, Supporting Information) reaches 0.077 S cm −1 at 600 °C and 0.35 S cm −1 at 700 °C (Figure 2d). In contrast, by using the reported D* values, [40] and the oxygen content and cell volume, [3] the calculated σ o of LSCF is ≈1.4 × 10 −4 S cm −1 at 594 °C, ≈8.7 × 10 −4 S cm −1 at 663 °C, and ≈7.9 × 10 −3 S cm −1 at 761 °C. Thus, σ o of BiSCF is at least two orders of magnitude higher than that of LSCF at similar temperatures, and is comparable to that of BSCF below 700 °C; above this temperature it is even higher (Figure 2d and Figure S11b, Supporting Information).
BiSCF also exhibits rapid oxygen surface exchange kinetics (Figure 2b). At a pO 2 of 0.21 bar, the oxygen surface exchange coefficient (k*) is ≈1.6 × 10 −7 and 3.2 × 10 −6 cm s −1 at 600 and 700 °C, respectively. Reported k* values strongly depend on measuring methods, sample form (thin film or ceramic), and exchange pO 2 . There are often significant discrepancies in reported k* values even for the same nominal composition, which can significantly depend upon the microstructure. [7,8] Nevertheless, overall, the k* values measured for BiSCF are slightly higher than those found for LSCF ceramic samples [40] but about one order of magnitude lower than those for BSCF ceramic samples measured using the same technique with similar exchange pO 2 . The E a for k* of BiSCF is 1.8 eV, which is close to that of LSCF (≈1.8 eV) measured above 600 °C. [40] The chemical composition of the outer surface of the BiSCF ceramic strongly depends on the thermal history of the samples. Low-energy ion scattering (LEIS) is a highly surface-sensitive technique to determine the composition of the outermost atomic layer of a material and has been employed to study ceramic, thin film, or powder samples. [42][43][44] LEIS analysis of the BiSCF ceramic using 5 keV Ne + ions shows the presence of all the cations after polishing, as expected for the starting composition (Figure 3a). Thermal treatment leads to a significant segregation of both Sr and Bi cations (as shown by an increase of the peak areas at ≈2150 and 3526 eV, respectively), with a concomitant decrease of the B cations surface coverage (indicated by the disappearance of the (Fe + Co) peak at ≈1370 eV after annealing at 650 °C). These results are consistent with previous LEIS measurements on ceramic LSCF where the outermost atomic layer of LSCF consists of A-site cations rather than B-site cations. [42,45] However, the ratios of the A-site cations in the outermost atomic layer of BiSCF differ from those in LSCF. [42,45] In LSCF, the peak corresponding to lanthanum vanishes after annealing at high temperatures (>600 °C) and the outermost atomic layer is shown to contain only Sr. [42,45] In BiSCF, the outermost atomic layer still contains more Bi and less Sr than the bulk (Figure 3b), consistent with the Bi surface enrichment also observed with LEIS in Bi-based pyrochlores. [29] The surface enrichment of Bi in BiSCF was also confirmed by X-ray photoelectron spectroscopy (XPS) data ( Figure S13, Supporting Information). XPS has a typical sampling depth of 3-10 nm and is less surface sensitive than LEIS. The sampling depth could be reduced by lowering the emission angles at which the electrons are collected, which enables detection of compositions closer to the surface. The intensity of Bi 4f peaks increases with lowering emission angles from 90° to 70° and 60° ( Figure S13 for the outermost atomic layer of BiSCF, the XPS results qualitatively support the Bi surface enrichment as revealed by LEIS. As Bi 3+ has a higher charge than Sr 2+ , the Bi surface enrichment presumably decreases the surface oxygen vacancy concentration which may in turn explain why BiSCF, despite its high oxygen mobility, exhibits a lower k* than BSCF at intermediate temperatures. Note that k* depends on not only intrinsic material parameters such as oxygen mobility and oxygen vacancy concentration [46] but also surface microstructure. [47][48][49] The approach of surface microstructure engineering could be employed to further improve the k* of BiSCF. In addition to the exceptional bulk oxide transport properties, BiSCF also exhibits high CO 2 tolerance (Figure 4). Rapid weight gain was observed for BSCF ( Figure 4a) due to reaction with CO 2 and formation of (Ba,Sr)CO 3 when heated in a CO 2containing atmosphere (70%CO 2 -20%O 2 -10%Ar) at 900 °C. The weight gain reached ≈10.0% after tests at 900 °C for 20 h. In sharp contrast, under the same test conditions the weight gain for BiSCF is only 0.35%. The XRD patterns (Figure 4b) reveal no observable SrCO 3 after test at 900 °C with 70% CO 2 for 20 h, whereas BSCF decomposed with significant formation of carbonates and CoO (Figure 4c).
The combination of high electrocatalytic performance and CO 2 tolerance of BiSCF was demonstrated through oxygen permeation measurements using 100% CO 2 sweep gas (Figure 4d). At high temperatures (950 °C) under He sweep gas, BiSCF exhibits a high oxygen permeation flux (JO 2 ) of 1.45 mL min −1 cm −2 , comparable to the 1.77 mL min −1 cm −2 obtained for BSCF. A sharp contrast between the CO 2 tolerance of BSCF and BiSCF was observed. Consistent with reports in the literature, [13,14] the JO 2 of BSCF at 875 °C dropped quickly to nearzero (<0.02 mL min −1 cm −2 ) when pure CO 2 sweep gas was introduced. For BiSCF, the JO 2 at 950 °C dropped from 1.45 to 0.87 mL min −1 cm −2 after ≈100 h and then became stable. The JO 2 recovered quickly to 1.57 mL min −1 cm −2 after the sweep gas was switched back to pure helium. The decrease of JO 2 under CO 2 sweep gas and the recovery of JO 2 by switching the sweep gas from CO 2 to helium have been commonly observed in the literature for many MIEC membranes including BSCF and LSCF. [6,13,14] The decrease of JO 2 under CO 2 sweep gas is attributed to the formation of carbonate phases if there is severe reaction between CO 2 and MIEC material, [13,14] or to reduced oxygen vacancy concentration associated with CO 2 adsorption on the membrane surface where the chemical reaction between CO 2 and MIEC material is less severe (e.g., in the case of LSCF). [6] After switching the sweep gas from CO 2 back to helium, JO 2 recovers due to the decomposition of carbonate phases and/or desorption of CO 2 from the membrane surface.
Adv. Mater. 2020, 32,1905200  The inset shows the expanded view of the CoO minority secondary phase, which is the same before and after the tests. c) The XRD patterns of BSCF after TGA tests in CO 2 showing significant decomposition of BSCF. The cubic phase (down-pointing triangle) has the BSCF-type structure but smaller lattice parameter. d) Oxygen permeation stability tests using 100% helium or CO 2 sweep gas for BiSCF and BSCF membranes. The membrane thickness is 1.0 mm. A BSCF membrane with the same thickness was also tested at 875 °C for direct comparison. At each temperature the JO 2 was initially tested in 100% helium. The sweep gas was then switched to 100% CO 2 for ≈260 h, before being finally changed back to 100% helium. The flow rate for the feed gas (21% O 2 , 79% N 2 ) and sweep gas is 100 mL min −1 . During these processes the roughness of the membrane surface can increase, which increases the surface exchange kinetics, and therefore JO 2 returns to be above the initial value. [50] After the tests at 950 °C, the temperature was decreased to 875 °C. The JO 2 at 875 °C with helium sweep gas was 1.25 mL min −1 cm −2 , which decreased to 0.37 mL min −1 cm −2 under CO 2 sweep gas, then recovered slowly to 0.58 mL min −1 cm −2 and became stable at this value during the 260 h test. The stable JO 2 values of BiSCF with pure CO 2 sweep gas at both 950 and 875 °C are higher than those of many other MIECs under similar test conditions. [2,14] After the tests at 875 °C, the temperature was increased to 950 °C. A 130 h test stability test in helium was performed. A stable JO 2 of 1.59 mL min −1 cm −2 was observed ( Figure S14, Supporting Information).
The high electrocatalytic performance of BiSCF was also apparent in its performance as a solid oxide fuel cell (SOFC) cathode. Symmetrical cathode/electrolyte/cathode cells were fabricated using BiSCF as cathode and LSGM (La 0.80 Sr 0.20 Ga 0.80 Mg 0.20 O 2.8 ) as electrolyte. The area specific resistance (ASR) of the BiSCF cathode is as low as 0.31 and 0.13 Ω cm 2 at 600 and 650 °C, respectively (Figure 5, Figure S16 and Table S6, Supporting Information). The ASR values of BiSCF cathodes are about one order of magnitude lower than the typical values of LSCF [51][52][53] and are comparable to those of BSCF without interfacial/microstructural optimization, [49] demonstrating the high electrocatalytic performance exhibited by BiSCF. There is scope for further improvement to the ASR of BiSCF through optimization of the microstructure. For example, the ASR of BSCF cathode is decreased when a porous samarium-doped ceria (SDC) interlayer is deposited between the BSCF cathode and the SDC electrolyte. [49] The E a of 1.3 eV is similar to that of BSCF and smaller than ≈1.7 eV of LSCF.
In summary, the substitution of the post-transition metal Bi 3+ cation onto the A-site of the perovskite structure in BSCF retains similar oxygen vacancy content and bulk oxide ion conductivity, while also significantly increasing the stability to CO 2 . As a result, BiSCF displays comparable electrocatalytic performance to BSCF in oxygen permeation membranes and SOFC cathodes, and, unlike BSCF, retains this performance in the presence of CO 2 . The high polarizability and diverse oxide coordination chemistry characteristic of the Bi 3+ "inert lone pair" outer shell electronic configuration confers comparable anion mobility while enhancing transition metal-oxide orbital overlap and electronic conductivity from the shorter BO bonding. Furthermore, the more electropositive character of Bi 3+ versus Ba 2+ and Sr 2+ decreases the basicity of the oxide anions in BiSCF below that in the alkaline-only A site BSCF where the ionic character of the AO bond is higher, and thus reduces reactivity with Lewis acidic CO 2 . This effect is amplified at the surface due to the locally increased Bi content. Together, these bulk and surface chemical effects allow BiSCF to largely retain its oxygen permeation performance in a pure CO 2 atmosphere, in stark contrast to BSCF. The work demonstrates an approach to achieve synergic tuning of the bulk and surface properties of a complex perovskite oxide mixed ionic-electronic conductor to couple high anion mobility with enhanced chemical stability.
The phase purity was examined using a Panalytical X'Pert Pro diffractometer with Co Kα 1 radiation. High-resolution time-of-flight (ToF) neutron powder diffraction patterns were collected from a powder sample contained in a thin-walled vanadium can at ambient temperature using the HRPD instrument at ISIS, UK. High-temperature ToF-NPD patterns were obtained from sintered pellets contained in a single-walled quartz gas-flow cell under flowing air, with data collected on warming in the range 500-800 °C using the Polaris instrument at ISIS, UK.
Ceramic microstructure and compositional analysis were performed using a Hitachi S-4800 field-emission scanning electron microscope (SEM) equipped with an Oxford INCA energy-dispersive X-ray spectroscopy (EDS) detector.
Cerimetric titration was used to determine the oxygen content in the as-prepared powder samples. The oxygen content at high temperatures Filled symbols indicate selected frequencies. Note that the ohmic impedance associated with the LSGM electrolyte has been subtracted from the spectra (an original impedance complex plane plot is shown in Figure S16, Supporting Information). The ASR values shown in the inset were normalized for the electrode area and divided by 2 to account for the symmetry of the cell. Comparison with typical ASR values of BSCF [49] and LSCF [51] is shown in the inset. and in air were determined using three methods. The first approach was to refine the high-temperature in situ NPD data. For the second approach, the samples were quenched from the desired temperature into liquid nitrogen in less than 0.5 s using a home-made rig ( Figure S1, Supporting Information). The oxygen content was then analyzed using the cerimetric titration method. The third approach was to first determine the oxygen content at 950 °C using the above quenching procedures. The oxygen contents at other desired temperatures were then obtained by measuring the weight difference between the desired temperature and 950 °C using TGA (TA Instruments SDT-Q600 TGA).
The CO 2 -tolerance of BiSCF and BSCF was evaluated by measuring the weight gain as a function of time using TGA under 70%CO 2 -20%O 2 -10%Ar at 900 °C.
The compatibility with common electrolyte materials (SDC, Sm 0.2 Ce 0.8 O 1.9 ; LSGM, La 0.80 Sr 0.20 Ga 0.80 Mg 0.20 O 2.8 ; and YSZ, 8 mol% yttria stabilized zirconia) was evaluated by annealing the mixtures of BiSCF and electrolyte materials at 850 and 950 °C for 3 h. The phase purity after annealing was examined using XRD.
A Linseis LSR -3 Linseis Seebeck Coefficient & Electric Resistivity Unit was used to measure dc conductivity and Seebeck coefficient in air from 100 to 800 °C using bar-shaped samples.
Oxygen isotope exchanges were conducted according to standard isotope exchange protocols in a custom-built exchange setup. The preannealing step in normal oxygen was approximately ten times as long as the subsequent 18 O exchange anneal (Table S1, Supporting Information). The oxygen pressure for both normal O 2 anneal and isotopically enriched O 2 anneals was 200 mbar. After the isotope exchange, the oxygen isotope distribution was determined by ToF-SIMS (ToF-SIMS V, Ion-ToF GmbH., Germany) analysis of polished cross sections of each sample. Typical diffusion profiles and fits for isotope diffusion profiles at 600 and 650 °C are shown in Figure S12 (Supporting Information). The chemical composition of the outermost atomic layer of the BiSCF ceramic was analyzed by LEIS (Qtac 100 Spectrometer, Ion-ToF GmbH, Germany).
Oxygen permeation measurements under an oxygen partial pressure gradient were performed in a Probostat (NorECs, Norway) measurement cell ( Figure S2, Supporting Information). Oxygen permeation fluxes were measured using a PerkinElmer Clarus 580 gas chromatograph (GC) equipped with a thermal conductivity detector.
Symmetrical BiSCF/LSGM/BiSCF cells were fabricated by screenprinting BiSCF ink on LSGM. Gold wire and gauze were attached to both sides of the symmetrical cell using gold paste. ASR measurements were performed at 550-750 °C.
Full experimental section and any associated references are available in the Supporting Information.

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.