Comprehensive Understanding of Cathodic and Anodic Polarization

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Energy, Environmental, and Catalysis Applications

A Comprehensive Understanding of Cathodic and Anodic Polarization Effects on Stability of Nanoscale Oxygen-Electrode for Reversible Solid Oxide Cells Sung Min Choi, Junsung Ahn, Ji-Won Son, Jong-Ho Lee, Byung-Kook Kim, Kyung Joong Yoon, and Ho-Il Ji ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b11874 • Publication Date (Web): 31 Oct 2018 Downloaded from http://pubs.acs.org on November 2, 2018

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ACS Applied Materials & Interfaces

A Comprehensive Understanding of Cathodic and Anodic Polarization Effects on Stability of Nanoscale Oxygen-Electrode for Reversible Solid Oxide Cells

Sung Min Choi,1 Junsung Ahn,1 Ji-Won Son,1,2 Jong-Ho Lee,1,2 Byung-Kook Kim,1 Kyung Joong Yoon,1 and Ho-Il Ji1,*

1

High-temperature Energy Materials Research Center, Korea Institute of Science and

Technology (KIST), Seoul 02792, Republic of Korea 2

Nanomaterials Science and Engineering, Korea University of Science and Technology

(UST), KIST Campus, Seoul 02792, Republic of Korea

*Corresponding author. E-mail address: [email protected],

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ABSTRACT: Degradation of oxygen-electrode in reversible solid oxide cells operating in both electrolysis and fuel-cell modes is critical issue that should be tackled. However, origins and mechanisms thereof have been diversely suggested mainly due to a difficulty in precise analysis of microstructural/compositional changes of porous electrode, which is a typical form in solid oxide cells. In this study, we investigate the degradation phenomena of oxygenelectrode under electrolysis and fuel-cell long-term operations for 540 h, respectively, using a geometrically well-defined, nanoscale La0.6Sr0.4Co0.2Fe0.8O3-δ (LSCF) dense film with a thickness of ~ 70 nm. Based on assessments of electrochemical properties and analyses of microstructural and compositional changes after long-term operations, we suggest consolidated degradation mechanisms of oxygen-electrode, including the phenomena of kinetic demixing/decomposition of LSCF, which is not readily observable in the typical porousstructured electrode. KEYWORDS: reversible solid oxide fuel cells, oxygen-electrode, polarization, degradation, thin film electrode

Solid oxide cells (SOCs) in a configuration of oxygen-electrode | oxygen-conducting electrolyte | fuel-electrode can operate in a reversible way as solid oxide fuel cells (SOFCs), converting a chemical fuel into electricity, and as solid oxide electrolysis cells (SOECs), converting electricity into the usable chemical fuel1-3. Whereas SOFCs have been widely investigated in the last two decades, it is not very long since SOECs were spotlighted due to increasing needs of on-site chemical energy storage systems4. Thus, most SOECs have been demonstrated using the already developed SOFCs with insufficient consideration on the electrode properties, although the transport of charge carriers and the electrochemical reactions at the electrodes in SOECs proceed in reverse directions of those in SOFCs5. As a result, 2 ACS Paragon Plus Environment

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specifically, the degradation issues on SOECs are being reported, which mainly stem from the oxygen-electrode (anode); delamination between the oxygen-electrode and the electrolyte, resulting in losses of electrical/ionic conductions6-7. Although Graves et al. reported that cyclic operations between fuel-cell and electrolysis modes with an interval of few hours effectively eliminate the degradation2, SOCs basically should be freely switched from fuel-cell to electrolysis modes or vice versa according to not only hour-by-hour, but even also season-byseason powder demands8. Therefore, it is currently recognized that modification/optimization of the oxygen-electrode are necessary to achieve securer long-term operations in the both electrolysis and fuel-cell modes. However, the systematic and consolidated understandings on degradation mechanism are still not clearly established due to the following reasons. Firstly, most studies addressing the stability and the degradation mechanism relevant to the oxygenelectrode in both of SOFC and SOEC modes have been performed for SOCs composed of porous-structured oxygen-electrodes, which inevitably possess not only an ambiguity on defining the geometric factors quantitatively but also a variability of microstructures depending on fabrication methods and process history. Furthermore, recent studies have applied nanostructured electrodes to enhance the performance via solution-based infiltration9. Thus the structural, compositional, and even spatial changes in the porous oxygen-electrodes cannot be readily assessed. Secondly, another factors such as constituents of SOCs (electrode and electrolyte materials, thickness of electrode, and ratio of each material in composite electrode) and operating conditions (temperature, atmospheric environment, and bias/current level), which are all closely correlated with the degradation phenomena, are not in agreement in most studies. The aforementioned complexity and diversity induce that the origins of degradation in the oxygen-electrode diversely suggested, not only in SOECs but also in SOFCs: for example, in regard to La1-xSrxCo1-yFeyO3- oxygen-electrode materials that are known to be superior to La1-xSrxMnO3- (LSM) electrodes in SOFCs as well as SOECs10, compositional fluctuations 3 ACS Paragon Plus Environment


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and re-crystallization6, microstructural changes, kinetic demixing with a formation of Co3O46, 11

, cation inter-diffusion between the oxygen-electrode and the electrolyte resulting in a

formation of SrCeO3 secondary phase12, and oxygen pressure build-up near an interface between the oxygen-electrode and the electrolyte13 have been suggested in SOECs. Regarding SOFCs, Sr segregation14, formation of SrCoOx secondary phase15, and cation inter-diffusion between the oxygen-electrode and the electrolyte16 have been reported. To avoid such limitations, here we investigate the degradation phenomena of the oxygen-electrode in electrolysis and fuel-cell operations, respectively, using half cells with a configuration of geometrically well-defined, nanoscale thin film La0.6Sr0.4Co0.2Fe0.8O3-δ (LSCF) deposited on Ce0.9Gd0.1O2-δ (GDC) electrolyte via pulsed laser deposition (PLD). Furthermore, to avoid any differences in fabrication process history, microstructure, and test conditions, two half cells were simultaneously prepared, stacked, and tested (Fig. 1a). One may think that the dense LSCF is probable to hinder the oxygen ion transport, and as a result, the degradation mechanism will be different with those of the actual porous-structured composite LSCF-GDC electrode. However, dense LSCF layer with a thickness of ~ 70 nm, i.e., having much shorter diffusion length than that of typical porous oxygen-electrodes (particle size is sub-microns), was used in this study, and an ionic conductivity of LSCF is known to be comparable with that of yttria-stabilized zirconia (YSZ)17-18. Consequentially, the nanoscale dense structure is rather beneficial to examine the structural and compositional variations after long-term operations compared to the porous structures. Furthermore, since LSCF exhibits high surface reaction activity associated with oxygen reduction/evolution reactions and high bulk diffusivity of oxygen ion19, the electrochemical reaction takes place mostly through two-phase-boundaries (interface between LSCF and gas phase). Thus, the configuration of thin film LSCF | (GDC electrolyte) can be considered as the simplified identical interfaces in the porous composite 4 ACS Paragon Plus Environment

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electrode such as LSCF | (GDC electrolyte) or LSCF | (GDC in composite electrode) interfaces. Using the thin and dense LSCF | GDC half cells, we analyze the electrochemical and microstructural/compositional changes of LSCF layers after long-term operations in electrolysis and fuel-cell modes, respectively. It will clarify the correlation between the electrochemical signals and the microstructural/compositional changes of the oxygenelectrodes, thereby enabling in operando diagnosis from the electrochemical analysis. Then we suggest the most probable degradation mechanisms and verify them in comparison with the degradation of SOECs comprised of the porous LSCF-GDC composite oxygen-electrode. As-prepared LSCF thin layer using PLD at 750°C on the polished polycrystalline GDC pellet was fully dense with a thickness of ~ 70 nm as shown in Fig. 1b. Electrochemical characteristics of the dense LSCF electrodes were investigated using quasi 4-probe configurations. For the oxygen-electrode under anodic polarization, i.e., in electrolysis mode (hereafter, A-electrode), Pt wires of (1) and (3) in Fig. 1a were used for working and counter electrodes, respectively, and (2) and (5) were used for reference electrodes. Similarly, for the oxygen-electrode under cathodic polarization, i.e., in fuel-cell mode (hereafter, C-electrode), Pt wires of (3) and (1) were working and counter electrodes, respectively, and (4) and (6) were reference electrodes. An equivalent circuit for a dense mixed-ionic-electronic conductor (MIEC) on an electrolyte has been suggested elsewhere20-21. Under a reasonable assumption that electronic resistances at the surface and in the bulk of LSCF are negligible due to a high electronic conductivity of LSCF, the equivalent circuit for this system can be simplified as s shown in Fig. 2a, where R ion is a reaction resistance associated with incorporation/release of

s oxygen ions at LSCF surface, Cion is an ionic capacitance at LSCF surface, R ion is a

 resistance of oxygen ion in bulk LSCF, Cchem is a chemical capacitance in bulk LSCF, R ion



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 is a resistance of oxygen ion transport at LSCF/GDC interface, Ceon is an electronic

capacitance at LSCF/GDC interface, and R 0 is a total resistance of GDC. Additionally, since the ionic conductivity of LSCF is known to be quite high, e.g., it is only ~2 times smaller than that of GDC at 650oC (LSCF: 0.009 S/cm17, GDC: 0.019 S/cm22), the resistances of ion 

transport in LSCF bulk ( R ion ) and through LSCF/GDC interface ( R ion ) can be neglected. It further reduces the equivalent circuit as shown in Fig. 2b. According to this model, impedance s spectrum will be displayed as one semi-circle corresponding to R ion  Ctot parallel circuit with

a high frequency offset corresponding to R 0 . Fig. 2c shows the experimental results of electrochemical impedance spectra for the both half cells under open circuit conditions at 650 and 600C in prior to long-term operations. Although they show little difference in values of the high-frequency intercept and the overall polarization resistance probably owing to a small difference in positions of reference electrodes, Pt wires of (5) and (6) in Fig. 1a, and in a contact area between the LSCF dense layer and the Pt mesh (here, precious metal pastes/patterns on the LSCF surface were not applied), the shapes of the impedance spectra were identical which means that two electrodes are in the same nature of the likely reaction pathway. The impedance spectra indeed well explained by the simplified equivalent circuit model in Fig. 2b, implying that the origins of degradation can be indirectly derived from the change behavior in the values s of R ion and R 0 .

The stacked half cells in Fig. 1a were simultaneously operated under constant current through the electrodes of (1) and (3) at 600°C. The current level of 2 mA/cm2 was determined based on I-V characteristics of the both A- and C-electrodes (Fig. S1a). The operating conditions (temperature, current density) and microstructural difference of nanoscale dense electrode are probable to change the degradation mechanism compared to that of conventional 6 ACS Paragon Plus Environment

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porous electrode. Thus, we carefully considered which operating conditions, i.e., temperature and current density, are adequate for the thin and dense electrode to exhibit the closest degradation behavior. First, the lower temperature of 600°C than the typical operating temperature of the actual SOCs (≥700°C) was chosen to minimize the effects from the instability of i) LSCF thin film itself at high temperature (deposition temperature is 750oC) and ii) LSCF material itself inducing such as Sr and/or Co segregations at the surface which typically observed above 700oC during long-term operation. Second, the lower current density of 2mA/cm2 was applied to minimize the effect of severer condition that possibly leads to the formations of pore and delamination due to the lower ionic conductivity and limited reaction pathways. If the ionic conduction through bulk path of LSCF were a rate limiting step, currentvoltage curve would show a diode-like asymmetric behavior. However, the actual currentvoltage curve for the thin film LSCF electrode (Fig. S1a) shows almost symmetric behavior, indicating that the current range up to ~4mA/cm2 does not correspond to the ion-blocking condition. The degradation behavior of the A-electrode was monitored by measuring voltage between the reference electrodes of (2) and (5) over time for 540 h, and the overall resistance of the half-cell continuously increased (Fig. S1b). After long-term operation, whereas the both s s of R 0 and R ion in the A-electrode increased (Fig. 2d), R 0 increased but R ion rather

decreased in the C-electrode (Fig. 2e). Since the bulk resistance of LSCF thin film is negligible due to the thickness of ~ 70 nm (8.3  10-4 ·cm2 based on the ionic conductivity of LSCF at 600°C17), the increase of R 0 is dominantly owing to the less contact area at LSCF/GDC s interface. The change of R ion can be postulated owing to the microstructural and/or the

compositional changes of LSCF layer. Additionally, small but non-negligible arcs were appeared at mid frequency regime (~ 103 Hz) in the both of fuel-cell and electrolysis modes as



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shown in insets of Fig. 2d-e, indicating that additional interfaces were possibly generated during long-term operations in the both cells (will be discussed below). Fig. 3 shows surface and cross-sectional images of the A- and C-electrodes after 540h long-term operations. As shown in Fig. 3a-b, LSCF layer in the cell under anodic polarization was locally delaminated (dark: LSCF, bright: GDC). Fig. 3c-d show the close-ups of LSCF surfaces. There existed blisters only at the surface of A-electrode, and cobalt segregations (as clearly shown in Fig. 3e-f, the very top layers) with a size of ~ 0.5 m were observed in the both cells, but their density was much higher in the C-electrode. Fig. 3g-h show cross-sectional microstructures in high magnification. Surprisingly, periodic pores with a size of ~ 40 nm were observed in the both cells; triangular pores near the electrolyte/electrode interface and circular pores near the electrode surface are formed in the A- and C-electrodes, respectively. Particularly, some circular pores in the C-electrode are expanded to the electrolyte/electrode interface. Additionally, as shown in Fig. 3i, the A-electrode is compositionally homogeneous (La : Sr : Co : Fe = 0.6 : 0.41 : 0.21 : 0.88 in atomic %) maintaining its original composition. Contrarily, based on a line along the circular pores in the C-electrode, two regions show a compositional difference (Fig. 3j). Whereas the composition of the lower region is in reasonable agreement with the initial composition (La : Sr : Co : Fe = 0.6 : 0.43 : 0.22 : 0.89), the upper region is Co-deficient and the ratio among the constituent cations (La : Sr : Co : Fe = 0.6 : 0.56 : 0.01 : 0.71) is deviated from the initial ratio (top layers are platinum deposited for SEM analysis in the both A- and C-electrodes, Fig. S2). To sum, the distinguished features induced by anodic and cathodic polarizations on LSCF oxygen-electrodes are i) delamination/blister originated from electrolyte/electrode interface, ii) degree of cobalt segregations, iii) location of internal pores, and iv) compositional variation. They can be correlated with the impedance changes after constant current operation for 540h discussed 8 ACS Paragon Plus Environment

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above (Fig. 2d-e). First, the increase of R 0 is owing to i) the less contact area between the electrolyte and the oxygen-electrode induced by the delamination (A-electrode) and by the pores existed at interface (A- and C-electrodes) and ii) the reduced conduction path resulted s from the internal circular pores (in C-electrode). Second, the increase of R ion in the A-

electrode is possibly originated from the less active area owing to the delamination/generation of blisters and the cobalt segregations at surface (poorly conducting Co3O4 phase6, 23). The s decrease of R ion in the C-electrode is somewhat unclear, since cobalt segregations are far

s severer than the A-electrode. However, Baumann et al., also reported the reduced R ion after

applying a strong cathodic polarization to La0.6Sr0.4Co0.8Fe0.2O3- thin film on YSZ, and suggested that severe changes of cation concentrations within the outermost surface layer is a main cause of the improvement3. Indeed, surface reaction rate constants of La1-xSrxFeO3- systems have been reported to be 2 ~ 6 times higher than that of LSCF24. Thus, it is most probable that Co-deficient LSCF layer at surface enhances the surface reaction rate, and in turn, s R ion is reduced. Third, the additional resistance corresponding to the arc at mid frequency in

impedance spectra for the both of fuel-cells after long-term operations (inset graphs in Fig. 2de) is known to be originated from LSCF/GDC boundary25. One probable reason resulting in 

the increase of interfacial resistance ( R ion ) is compositional changes at the electrolyte/electrode interface owing to kinetic demixing of LSCF26-27 or cation inter-diffusion between the electrolyte and the electrode16. However, elemental line scan/mapping results (Fig. S3-4) indicate that there are no observable compositional changes. Thus, the pores at interface in both samples (Fig. 3g-h) are highly probable to be the main source of the relevant resistance by hindering facile oxygen ion transport through the interface.



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Based on the observations and analyses thereof, the oxygen-electrode degradation mechanisms under anodic (electrolysis) and cathodic (fuel-cell) polarizations are suggested (Fig. 4). Regarding anodic polarization, it has been suggested that the locally higher oxygen pressure near electrolyte/oxygen-electrode interface leads to the oxygen-electrode delamination13. Thus, uses of i) oxygen-electrodes exhibiting high ionic conductivity and high activity in oxygen evolution reaction and ii) electrolytes with relatively high electronic conductivity are beneficial to lower the oxygen pressure build-up at interface. In this regard, a combination of LSCF and GDC is expected to be better than a combination of LSM and YSZ to prevent the delamination. However, the geometry of flat film of LSCF, which is used in this study, has low surface area compared to the porous electrode, i.e., oxygen evolution reaction is not facile enough, and in turn, substantially higher propensity to be delaminated. As a result, pores are generated from GDC/LSCF interface, then grow toward the surface of LSCF as the electrolysis operation is prolonged. Finally, blisters and delamination take place and the performance is degraded (Fig. 4a). Regarding cathodic polarization, an oxygen activity of internal LSCF decreases as the fuel-cell operation proceeds due to an insufficient rate of oxygen electro-reduction at surface of oxygen-electrode and/or a relatively lower ion conductivity of LSCF than that of GDC. When the oxygen chemical potential gradient between internal LSCF and GDC near at interface exceeds a certain point, cations of LSCF at subsurface are prone to move towards free surface (kinetic demixing/decomposition). Indeed, such processes driven by identical origin, i.e., high oxygen chemical potential gradient, have also been extensively observed in LSCF oxygen-separation membranes; compositional gradient/formation of secondary phases near/at an oxygen-rich side26, 28. Thus, the formation of pores near LSCF surface is highly probable to be originated from the kinetic demixing/decomposition processes. Furthermore, as B. Wang et al., reported that cobalt cation diffusivity is the highest among the constituent cations in LSCF29 and X. Wang et al., reported 10 ACS Paragon Plus Environment

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that A-site-rich and B-site-rich LSCFs tend to form Sr and Co segregations on the surface, respectively30, Co segregations at the surface and a consequent compositional gradient in the region above the line along pores in the C-electrode (Fig.3d) are not surprising. In the same manner, less Co segregation at surface of the A-electrode seems to be obvious, since the oxygen chemical potential gradient is applied in the reverse direction. Degradation phenomena, which was observed in an actual porous oxygen-electrode in a form of LSCF-GDC composite after electrolysis operation for ~1300 h under current density of 2A/cm2 at 700oC, can be explained by the suggested mechanisms above. The interface between porous LSCF-GDC oxygen-electrode and GDC electrolyte was partly delaminated after prolonged electrolysis operation (Fig. S5). Particularly, in a region where the adhesion is maintained but near the delaminated region (Fig. 5b), LSCF phases (dark color) are more distributed than the initial state at interface (Fig. 5a). As a result of elemental mapping for LSCF near interface (Fig. 5c), concentrations of Co and Sr (in region 3, La : Sr : Co : Fe = 0.6 : 0.92 : 0.62 : 0.74) are higher than those of original LSCF composition. This result indicates the phenomena that we observed above such as the migration of cations in LSCF towards interface driven by oxygen chemical potential gradient (reverse direction of that in C-electrode, Fig. 4b) and the kinetic demixing/decomposition resulting in compositional variance happen simultaneously. Finally, the interface is delaminated by the identical mechanism suggested in Fig. 4a, because the less ion-conductive phases compared to the electrolyte covers the interface more and more over time. In summary, we investigated the degradation phenomena in nanoscale LSCF thin film in electrolysis and fuel-cell operations. Owing to the benefit of well-defined geometry of LSCF thin film, the microstructural/compositional changes in LSCF films were successfully analyzed in nanoscale, and the correlation between the components of electrochemical impedance and 11 ACS Paragon Plus Environment


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the major origins resulting in degradations was clarified. It is expected to enable in operando diagnosis of solid oxide cells during long-term operations simply from the electrochemical analysis. Furthermore, we suggested the most probable degradation mechanisms, and importantly, it was newly suggested that kinetic demixing/decomposition of LSCF, which is not readily observable in the typical porous-structured electrode, are highly probable to affect the both fuel-cell and electrolysis long-term degradations. The suggested mechanisms were able to further explain the degradation phenomena occurring in LSCF-GDC porous oxygenelectrode under electrolysis operation.

ASSOCIATED CONTENT Supporting Information: The Supporting Information is available free of charge on the ACS Publications website. Experimental

section;

additional

electrochemical,

microstructural,

compositional

characterizations on dense film electrode and bulk porous electrode (Figure S1-S5)

ACKNOWLEDGMENT This research was supported by Technology Development Program to Solve Climate Changes through the National Research Foundation (NRF) of Korea funded by the Ministry of Science, ICT (NRF-2016M1A2A2940148) and the institutional research program of Korea Institute of Science and Technology (KIST).

REFERENCES


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(20) Jamnik, J.; Maier, J. Generalised Equivalent Circuits for Mass and Charge Transport: Chemical Capacitance and Its Implications. Physical Chemistry Chemical Physics 2001, 3, 1668-1678. (21) Usiskin, R. E.; Maruyama, S.; Kucharczyk, C. J.; Takeuchi, I.; Haile, S. M. Probing the Reaction Pathway in (La0.8Sr0.2)0.95MnO3+δ Using Libraries of Thin Film Microelectrodes. Journal of Materials Chemistry A 2015, 3, 19330-19345. (22) Kharton, V.; Figueiredo, F.; Navarro, L.; Naumovich, E.; Kovalevsky, A.; Yaremchenko, A.; Viskup, A. P.; Carneiro, A.; Marques, F.; Frade, J. Ceria-Based Materials for Solid Oxide Fuel Cells. Journal of Materials Science 2001, 36, 1105-1117. (23) Sharma, V. I.; Yildiz, B. Degradation Mechanism in La0.8Sr0.2CoO3 as Contact Layer on the Solid Oxide Electrolysis Cell Anode. Journal of the Electrochemical Society 2010, 157, B441-B448. (24) Armstrong, E. N.; Duncan, K. L.; Wachsman, E. D. Effect of A and B-Site Cations on Surface Exchange Coefficient for ABO3 Perovskite Materials. Physical Chemistry Chemical Physics 2013, 15, 2298-2308. (25) Baumann, F. S.; Fleig, J.; Habermeier, H.-U.; Maier, J. Impedance Spectroscopic Study on Well-Defined (La, Sr)(Co, Fe)O3-δ Model Electrodes. Solid State Ionics 2006, 177, 1071-1081. (26) Lein, H. L.; Wiik, K.; Grande, T. Kinetic Demixing and Decomposition of Oxygen Permeable Membranes. Solid State Ionics 2006, 177, 1587-1590. (27) Oh, M.; Unemoto, A.; Amezawa, K.; Kawada, T. Material Stability and Cation Transport of La0.6Sr0.4Co0.2Fe0.8O3-δ in SOFC Cathodic Conditions. Ecs Transactions 2011, 35, 2249-2253. (28) Wang, B.; Zydorczak, B.; Wu, Z.-T.; Li, K. Stabilities of La0.6Sr0.4Co0.2Fe0.8O3-δ Oxygen Separation Membranes—Effects of Kinetic Demixing/Decomposition and Impurity Segregation. Journal of Membrane Science 2009, 344, 101-106. (29) Wang, B.; Zydorczak, B.; Poulidi, D.; Metcalfe, I.; Li, K. A Further Investigation of the Kinetic Demixing/Decomposition of La0.6Sr0.4Co0.2Fe0.8O3-δ Oxygen Separation Membranes. Journal of Membrane Science 2011, 369, 526-535. (30) Wang, X.; Miyazaki, T.; Yashiro, K.; Hashimoto, S.; Kawada, T. The Origin of Instability of Lanthanum Strontium Cobalt Ferrite (La–Sr–Co–Fe–O; LSCF) under Oxygen Potential Gradient. ECS Transactions 2017, 75, 1-9.


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Figure 1. a) Schematic illustration of stacked two half cells for simultaneous electrolysis and fuel-cell operations. b) Cross-sectional transmission electron microscopy (TEM) image of asprepared LSCF thin film on GDC. Inset is scanning electron microscopy (SEM) image of LSCF surface.



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Figure 2. a) Equivalent circuit for mixed ionic-electronic conductor. b) Reduced equivalent circuit from a) for LSCF|GDC configuration. c-e) Impedance spectra (Nyquist plot) under open circuit condition for two half cells at 650 and 600oC in air before constant current long-term operation (c), before/after electrolysis operation for 540h at 600oC (d), and before/after fuelcell operation for 540 h at 600oC (e).


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Figure 3. Effects of long-term operation on structure of LSCF thin film. a-b) Surface views of A- and C-electrodes. c-d) surface SEM images. e-f) Energy dispersive spectrometer (EDS) elemental mapping at electrode cross-sections. g-h) Cross-sectional TEM images. i-j) EDS elemental mapping in high-magnifications of (e-f). La:Sr:Co:Fe ratios are calculated by averaging concentrations of each element at marked regions. 17 ACS Paragon Plus Environment


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Figure 4. Schematic illustration of degradation mechanisms via anodic polarization (a) and cathodic polarization (b).


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Figure 5. a-b) cross-sectional SEM images of porous LSCF-GDC composite electrode before (a) and after long-term electrolysis operation for ~1300h (b). c) Cross-sectional TEM image at porous electrode/GDC interface (left) and EDS elemental mapping of La, Sr, Co, and Fe. At region 1, 2, and 3, the ratios of La : Sr : Co : Fe are 0.60 : 0.30 : 0.16 : 0.74 (region 1), 0.60 : 0.35 : 0.17 : 0.71 (region 2), and 0.60 : 0.92 : 0.62 : 0.74 (region 3), respectively.



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TOC GRAPHICS

Anodic Polarization

Cathodic Polarization cation

O2

LSCF

GDC

O2

LSCF

GDC

O2-

100 nm

O2-

100 nm


Comprehensive Understanding of Cathodic and Anodic Polarization

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