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Over a >1100 h testing at 800 °C, the overcoated LSCF-GDC cathode exhibited respective polarization and ohmic area-specific-resistances 3 and 1.5 tim...
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Atomic Layer Deposition Functionalized Composite SOFC Cathode La0.6Sr0.4Fe0.8Co0.2O3‑δ -Gd0.2Ce0.8O1.9: Enhanced Long-Term Stability Yunhui Gong,† Rajankumar L. Patel,‡ Xinhua Liang,‡ Diego Palacio,§ Xueyan Song,§ John B. Goodenough,⊥ and Kevin Huang*,† †

Department of Mechanical Engineering, University of South Carolina, Columbia, South Carolina 29201, United States Department of Chemical and Biochemical Engineering, Missouri University of Science and Technology, Rolla, Missouri 65409, United States § Department of Mechanical & Aerospace Engineering, West Virginia University, Morgantown, West Virginia 26506, United States ⊥ Texas Materials Institute, The University of Texas at Austin, Austin, Texas 78713, United States ‡

ABSTRACT: We report that the long-term stability of a conventional mixed oxide-ion and electron conducting solid oxide fuel cell cathode, La0.6Sr0.4Fe0.8Co0.2O3−δ-Gd0.2Ce0.8O1.9 (LSCF-GDC) composite, can be markedly improved by functionalizing its surfaces with a conformal layer of nanoscale ZrO2 films through atomic layer deposition (ALD). Over a >1100 h testing at 800 °C, the overcoated LSCF-GDC cathode exhibited respective polarization and ohmic area-specific-resistances 3 and 1.5 times lower than the pristine sample, whereas the pristine LSCF-GDC cathode degraded at a rate 4 times faster than the overcoated one. The multifunctionality of porosity, mixed conductivity, and suppressed Srenrichment enabled by the nanoscaled ALD-ZrO2 overcoats are attributed to the performance retention observed for the overcoated cathode.

KEYWORDS: atomic layer deposition, solid oxide fuel cell, cathode, stability

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INTRODUCTION The mixed oxide-ion and electron conducting perovskite oxide La 1−x Sr x Fe 1−y Co y O 3−δ (LSCF) and its mixture with GdxCe1−xO2 (GDC) is a promising cathode for intermediate temperature solid oxide fuel cells (IT-SOFCs) because of its high intrinsic electrocatalytic activity to the oxygen reduction reaction (ORR) and manageable thermal expansion coefficient. However, the utmost challenge to the deployment of the LSCF-GDC cathode into commercial IT-SOFCs is the performance instability over a prolonged operation.1−5 One source of the degradation originates from Cr-poisoning where high-valence gaseous Cr-species result from a long-term exposure of Cr-containing oxidation-resistant metal interconnects to high temperature and a wet oxygen environment; these species are electrochemically reduced (condensed) into solid Cr2O3(s) at three-phase as well as two-phase boundaries, blocking ORR reactive sites.6−8 Coating of dense and conductive ceramic layers on the surface of metal interconnects as a means of mitigating the Cr-poisoning effect has proven rather effective.9−11 Another source of performance degradation arises from the LSCF itself. It is commonly observed that a LSCF or LSCFGDC IT-cathode can still experience pronounced performance degradation even in laboratory tests absent of metal interconnects.1−5 A wide range of surface analysis techniques, either in situ or ex-situ, have provided compelling evidence linking the performance decay in LSCF to the surface © 2013 American Chemical Society

segregation of Sr species in the form of SrO(s) or SrCO3,12−16 a chemical process that is sensitive to temperature, partial pressure of oxygen (pO 2) and electrochemical potential. 10,12−16 Since the ORR activity of an AMO 3 perovskite cathode is critically determined by the atomic structure (e.g., electronic configuration) and composition (e.g., cation concentration and oxygen nonstoichiometry) on the surface,17−21 a complete coverage of the passive and insulating SrO(s) layer over the cathode surface would easily block the ORR-active sites for effective charge-transfer. Compared to recent progress made in understanding the surface Srenrichment process, a means of alleviating such a detrimental phenomenon is largely lacking. We herein report that the conventional LSCF-GDC composite cathode can retain its electrochemical performance with improved long-term stability at 800 °C for >1100 h by functionalizing its surfaces with a conformal layer of nanoscaled ZrO2 films through atomic layer deposition (ALD). ALD is an emerging cutting-edge thin-film technology established by the unique binary-sequence and self-limiting reaction chemistry capable of yielding highly conformal and uniform nanoscaled films on the surfaces of almost any open bulk geometries.22−30 Received: June 30, 2013 Revised: October 10, 2013 Published: October 16, 2013 4224

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Figure 1. TEM diffraction contrast images showing (a) the porous LSCF backbone structure overcoated with a uniform and conformal layer of ZrO2 nanofilms; (b) a magnified view of image a showing a clear view of conformal ZrO2 coating; (c) and (d) HRTEM images showing 1−2 layers of randomly orientated ZrO2 grains (marked by the dotted lines); (e) indexing for the diffraction spots from the LSCF backbone; (f) indexing of the (101) rings from ZrO2 grains, showing cubic and tetragonal structures for LSCF and ZrO2, respectively. The sample was thermally treated at 600− 800 °C for 4 h.

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(600 s)→ evacuation (30 s). Two levels of ALD cycles, namely 30 and 60, were tried to make ZrO2-overcoats with different thicknesses. Electrochemical Impedance Spectroscopy Analysis. We employed electrochemical impedance spectroscopy (EIS) as the technique to evaluate the ohmic and ORR-related polarization ASRs of a cathode. An electrochemical workstation (IM6, Zahner) was used to conduct all the EIS measurements under an open circuit condition from 600 to 800 °C with 50 °C interval and long-term stability test at 800 °C. The frequency was swept from 1 × 105 to 0.05 Hz with an AC stimulus amplitude of 10 mV. Silver paste (C8829, Heraeus) and silver mesh served as the current collector for all the measurements. To minimize the variability, EIS spectra of the pristine and ALDZrO2 overcoated cathodes were simultaneously measured in a side-byside configuration in the same test fixture for the long-term test. The long-term test was conducted at 800 °C. The selection of 800 °C was intentional: to accelerate cathode degradation by which the effect of ALD-ZrO2 can be best evaluated within a shorter time frame. The results are deemed applicable to lower temperatures. We have recently demonstrated the remarkable effectiveness of ALD-ZrO2 on stabilizing nanostructured Sr-doped LaCoO3 cathode operated at 700 °C in a separate study.35 The short-term tests were performed as a function of temperature ranging from 600 to 800 °C, typically lasting approximately 4 h. For all the tests, air was constantly supplied to the samples at a flow rate of 50 sccm. TEM Examination. A high-resolution transmission electron microscope (HRTEM) equipped with energy-dispersive X-ray spectroscopy (EDS) was used to examine the structure and chemistry from micrometer to atomic levels. TEM samples were prepared by mechanical polishing and ion milling in a liquid-nitrogen cooled holder. Electron diffraction, diffraction contrast, high-resolution TEM imaging, and chemical EDS-analysis were performed with a JEM-2100 operated at 200 kV.

EXPERIMENTAL SECTION

Preparation of LSCF-GDC Composite Cathode in Symmetrical Impedance Cells. A symmetric cell configuration of cathode/LSGM/cathode was employed to evaluate area specific resistances (ASRs) associated with the cathode and electrolyte. The LSGM (La0.80Sr0.20Ga0.83Mg0.17O3−δ) electrolyte membranes were fabricated by a standard tape-casting method as previously described in ref [31, 32], the thickness and diameter of which were 200 μm and 15 mm, respectively. The rational to choose LSGM as the electrolyte for the cathode characterization is 2-fold: (1) LSGM would not react with LSCF during fabrication and operation, thereby avoiding the use of a GDC buffer layer if an YSZ electrolyte is otherwise employed; (2) LSGM has a higher ionic conductivity, by which the charge-transfer process, the focus of this study, can be maximally displayed over an impedance spectrum of a fixed time domain. To make the LSCF-GDC cathode, a commercial paste purchased from Fuel Cell Materials (LSCFGDC-1, sku: 232202) was screenprinted on both sides of the LSGM electrolyte membrane, followed by firing at 1100 °C for 1 h. The resultant cathode possesses a porous microstructure with an effective surface area of 0.75 cm2 and thickness of 30 μm. Atomic Layer Deposition of ZrO2 Overcoats on LSCF-GDC Composite Cathode. A flow-type reactor was used to deposit nanoscaled ZrO2 films directly on the exposed surface of cathode. Our reactor system is very similar to the one described previously.33 Tetrakis(dimethylamido)zirconium(IV) (electronic grade, ≥99.99%, Sigma-Aldrich) and deionized water were used as the precursors.34 In a typical run, a number of LSCF-GDC/LSGM/LSCF-GDC symmetrical cells were placed on a metal-mesh basket in the middle of the reactor. The reaction temperature was chosen as 250 °C. The Zr precursor was first vaporized in a bubbler at 80 °C before being delivered into the reactor for the Zr dosing. Room temperature water-vapor was used for oxidizer dosing. Before the reaction, the samples were first degassed at 250 °C for 5 h. The first precursor was then introduced for sufficient time to ensure that all exposed surfaces in the samples were saturated. After each dosing, a pure N2 stream at a flow rate of 15 sccm was then introduced to purge out the unreacted precursor and any byproducts formed during the preceding reaction. The system was then pumped down to 50 mTorr again prior to dosing the next precursor. A typical coating cycle followed the sequence: Zr precursor dose (60 s)→ N2 purge (600 s)→ evacuation (30 s)→ water dose (60 s)→ N2 purge

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RESULTS AND DISCUSSION Microstructure of the LSCF-GDC Cathode Overcoated with ALD-ZrO2. The TEM diffraction contrast images of the ZrO2 overcoats on the porous LSCF-GDC cathode backbone with 30 ALD cycles and a thermal exposure to 600−800 °C for approximately 4 h are shown in images a and b in Figure 1. As expected, these overcoats are uniform in thickness and 4225

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Figure 2. HRTEM images of the ALD-derived ZrO2 nanoscale overcoats. (a) Cross-sectional view showing through-thickness pores; (b−f) plan views of the stand-alone nanoscaled ZrO2 films showing grains/grain-boundaries and the preferential locations of pores formed at grain-boundaries.

in tetragonal phase. The second evidence is provided by a study that shows the formation of tetragonal and cubic ZrO2 with nanoparticle precursors.39,40 It is believed that the high surface energy owned by the nanoscaled grains plays a crucial role in the phase transformation. The third evidence is furnished by images e and f in Figure 1 of this study, in which a tetragonal (a = b = 3.598 Å, c = 5.152 Å) and cubic (a = 3.86 Å) structure are revealed for ZrO2 nanograins and backbone LSCF grains, respectively, by a selected area electron diffraction (SAED). By analyzing the electron diffraction rings from the ZrO2, the possibility of existence of monoclinic and cubic ZrO2 phase from the overcoat was excluded. Overall, the combined O2−-migration favorable crystal structure and mixed oxide ion and electronic conduction discovered in ALD-ZrO2 overcoats provide essential elements for a fast ORR catalysis. The counter-diffusion of Zr into the LSCF backbone, on the other hand, suppresses the surface Srenrichment through a point-defect charge-compensation mechanism to be discussed in the following. Another important microstructural feature of the ZrO2 overcoats is the existence of porosity. These pores are marked in Figure 2a, where the cross-sectional view of the film is shown. Identification of porosity is based on the observation of disrupted diffraction fringes from the lattice by the noncrystalline phases where the resin-filled pores are located. Those pores are deemed amorphous since there are no crystal fringes appeared in the marked region. An EDS analysis shown in Figure 2b inset on these regions indicates a significant amount of carbon signal (a main element of resin), which is undetectable on either LSCF grains or ZrO2 overcoats. The pores are seen in the range of a few nanometers as continuous channels through the thickness of the film. Observation of nanoscaled porosity in ALD derived thin films has been well documented in the literature, and its origin is linked to the dehydration and removal of carbon residues in the original assynthesized films during thermal exposure.30,41 Because the polycrystalline films are only 1−2 grains thick, formation of these nanoscaled porous channels would be a low-energy process. Similar to images c and d in Figure 1, the streaks at the LSCF/ZrO2 interface shown in Figure 2a are also related to the moiré fringes.

completely conformal around the grains of LSCF and GDC. The thickness is estimated to be ∼20 nm, inferring a rather fast deposition rate at ∼2/3 nm per each ALD-cycle. The crosssectional views of the ZrO2 overcoats were further imaged with HRTEM and are shown in images c and d in Figure 1. The films are seen to consist of 1−2 layers of randomly oriented grains (distinct grains are marked by yellow dotted lines) sized in 10−15 nm, suggesting a polycrystalline nature of the films. The discrete darker spots in Figure 1a of diffraction contrast image are caused by the nanograins with different crystal orientations in the polycrystalline films. The random crystal orientation is also consistent with the electron diffraction rings, Figure 1f, obtained from the ZrO2 overcoats. It is to be noted that the fringes at the interface between LSCF and ZrO2 overcoats as shown in Figure 1c, d are the moiré fringes36 originated from the overlapping of the LSCF with ZrO2 crystals. Because of the granular features of the LSCF crystals, when the electron beam is aligned with the low index zone axis of LSCF for taking HRTEM images, the interface between the LSCF and ZrO2 are not at the edge-on condition with respect to the electron beam. As a consequence, there is overlapping between ZrO2 and LSCF grains along the electron beam direction, thus causing the appearance of the moiré fringes. Local EDS-analysis found no trace of Sr and La, but approximately 6.0 at% Fe and 1.7 at% Co in the ZrO2 overcoats and 3.0 at% Zr in LSCF in the region close to the LSCF/ZrO2 interface. These mutual interdiffusions of Fe/Co and Zr are expected to progress with time until thermodynamic equilibrium is reached. It is also worth pointing out that neither SrZrO3 nor La2Zr2O7 is found anywhere within the layers including interfaces, consistent with thermodynamic prediction. The observation of Fe/Co in ZrO2 and Zr in LSCF has two implications: introduction of mixed conductivity and stabilization of high-symmetry crystal structures to lower temperatures. While the introduction of mixed conductivity by the mixedvalent Fe and Co cations is straightforward, attainment of highsymmetry structure in ALD-ZrO2 requires further experimental proofs. The first evidence is given by a prior study on ALDZrO2 doped with Fe,37,38 in which ZrO2 was found to crystallize 4226

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Figure 3. Comparison of EIS spectra among the pristine, 30-cycle, and 60-cycle ZrO2-overcoated LSCF-GDC cathodes in the “fresh” state. The EIS spectra were measured at 800 °C and in a flowing air.

Figure 4. Arrhenius plots of (a) RO and (b) RP(HF) and RP(LF) for both pristine and ZrO2-overcoated LSCF-GDC cathodes.

The plan-views shown in Figure 2b−f of the film alone, on the other hand, further reveal the polycrystalline nature of the films with pores preferentially located at the grain boundaries. These through-thickness nanoscaled porous channels formed during the thermal exposure make O2 accessible to the embedded LSCF for the ORR. This feature, along with the mixed oxide ion and electron conduction in high-symmetry structure, is deemed essential to the retention of ORR activity observed in the ALD-ZrO2 overcoated LSCF-GDC. Electrochemical Behaviors of the LSCF-GDC Cathode Overcoated with ALD-ZrO2. The effects of ALD cycles, temperature and time on electrochemical properties of the pristine and ZrO2-overcoated LSCF-GDC cathodes are studied by electrochemical impedance spectroscopy (EIS). Two important pieces of information can be obtained from an EIS spectrum measured: (1) ohmic area-specific resistance (RO) represented by the highest-frequency intercepts with the Z′-axis and (2) total polarization area-specific resistance (RP)

represented by the length on the Z′-axis between the highestand lowest-frequency intercepts. The ALD-cycle effect on EIS spectra is shown in Figure 3. An equivalent electrical circuit shown in the inset is established to simulate the EIS spectra measured. The basic shapes of EIS spectra are quite similar for both the pristine and overcoated samples, and the associated electrode process can be modeled by two R-CPE parallel subcircuits with sufficient accuracy (CPE: constant phase element). According to the literature,42,43 the high-frequency arc usually corresponds to the charge-transfer process, whereas the low-frequency arc relates to the dissociative adsorption of oxygen molecules. The areaspecific resistances of these two processes are denoted as RP(HF) and RP(LF) in this study, respectively. Figure 3 unarguably shows that the pristine sample exhibits the lowest RO and RP (= RP(HF) + RP(LF)), followed by 30cycle and 60-cycle coated samples, respectively. This is understandable in that the initially “fresh”, dense and nonconductive ZrO2 films present a barrier to electrons flow, 4227

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Figure 5. (a) Temporal evolution of EIS spectra of pristine and ZrO2-overcoated (30 cycle) LSCF-GDC cathodes as a function of time; Variations of (b) RO, (c) RP(HF), and (d) RP(LF) with selected time intervals.

Figure 6. Comparison of long-term stability of the pristine LSCF-GDC cathode with the ZrO2-overcoated one at 800 °C; (a) RO and (b) RP.

O2-diffusion and ORR charge-transfer. As is shown later, the RO and RP of these samples also show “aging” effect. The evolution of RP with time provides important hints to the understanding of the rate-limiting step in the overall ORR process of the ALDZrO2 overcoated LSCF-GDC. The effect of temperature on RO and RP (RP(HF) and RP(LF)) is further compared in Figure 4 in Arrhenius format for the pristine and overcoated LSCF-GDC cathodes. Note that the samples measured were also in “fresh” state. As expected, both RO and RP of the pristine sample are smaller than those of the overcoated ones. The Arrhenius plots in Figure 4a yield activation energies Ea in the range of 0.58−0.66 eV for RO of the pristine and overcoated samples. These values are comparable to that of LSGM electrolyte reported previously,31 which suggests that oxide-ion conduction in the LSGM electrolyte is a major source for the ohmic resistance. On the other hand, Figure 4b shows that the dissociative-adsorption related low-frequency RP(LF) has an Ea twice that of electrontransfer related high-frequency RP(HF) for both samples. The range of Ea observed is consistent with those reported in the

literature where the dissociative adsorption generally needs more energy to activate than the electron transfer.44−46 Further comparison between the two samples also suggests a similar Ea for the same electrode process, e.g., either dissociative adsorption or electrons transfer. Therefore, the higher RP exhibited by the “fresh” overcoated sample is likely attributed to a lowered surface concentration of adsorbed active oxygen species by the presence of the initially insulating ZrO2overcoats. The time “aging” effect on RO and RP of the pristine and ZrO2-overcoated cathodes reveals the most interesting results. Figure 5a shows the temporal evolution of EIS spectra for the pristine and overcoated cathodes at several selected time intervals. In general, during this period, the pristine sample exhibits a persistent increase in total RP (= RP(HF) + RP(LF)) and RO (more pronounced after 100 h) with time, whereas the overcoated sample shows a much smaller degree of increase in RP, and even a surprising decrease in RO for the first 100 h, see Figure 5b. The separation of RP into RP(HF) and RP(LF) from Figure 5a makes it even more clear that a large portion of RP 4228

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reduction for the overcoated sample originates from RP(LF). In other words, for a particular sample, RP(LF) has a more pronounced “aging” effect than RP(HF), suggesting that the dissociative adsorption of oxygen molecules is mostly affected by the time on temperature. This finding is consistent with previous studies where the Sr-enrichment is identified as the root cause for the deactivation of oxygen dissociative adsorption on the surface of a LSCF cathode.12−16 The relative rates of change in total RP (= RP(HF) + RP(LF)) and RO are further compared in Figure 6 for the pristine and ZrO2-overcoated samples over a continuous 1100 h time-scale. It is found that RP value of the pristine sample at the 1100 h marker has grown by nearly 3 times higher than the overcoated sample. Similarly, RO of the pristine sample at the 1100-h marker has increased by nearly 1.5 times higher than the overcoated one, whereas RO of the overcoated sample exhibits a surprising initial decrease, followed by a plateau afterward. Such simultaneous and significant increases in RP and RO have resulted in the final degradation rate (or the rate of increase in RP + RO) of the pristine cathode 4 times faster than the overcoated one. Overall, the ALD-ZrO2 overcoats show an excellent ability to stabilize the ASRs of the original LSCFGDC cathode. It should be noted that neither the absolute ASR values nor degradation rate obtained from the commercial LSCF-GDC cathode of this study were impressive when compared to the published data; typical values of RP and degradation rate for the same cathode are in the order of 0.1 Ω cm2 and ∼2% (within 200 h), respectively, at 750−800 °C.44−47 It is impossible to speculate the reason for the poor performance observed because of a lack of information on the exact makeup and fabrication condition of this commercial product. However, because the present study is comparative in nature, the results are unquestionably applicable to other better performed cathodes just as reported in our recent work.35 Defect-Chemistry Perspective. In an effort to understand why and how the ALD-derived nanoscale ZrO2 overcoats can improve the thermal stability of the original LSCF-GDC cathode, the following defect-chemistry model is proposed. It is known that the retention of oxygen vacancies,V•• O , in air at the cathode requires operation on a redox couple that is near or pinned at the top of the O-2p band.48 This condition also provides a large enough admixture of M-3d and O-2p orbitals in the redox couple so that the (180°-ϕ) M-O-M d-orbital interactions are strong enough to give itinerant holes in the M4+/M3+ redox couple created by the as-prepared Sr/La. From a defect-chemistry perspective, there is a surface reaction (in Kröger-Vink notation) 1 × × • O2 + V •• O + 2MM = OO + 2M M 2

2MM×

The MIEC condition optimizes the ORR activity of a SOFC cathode. The availability of abundant electron holes and oxygen vacancies in LSCF under oxidizing atmospheres also becomes / the driving force to form SrO(s) by oxidizing SrLa , a substitutional point defect created by replacing La with Sr, accompanied by the reduction of M•MtoMM× (M = Fe and Co) with free electrons transferred from the σ*-band of M-3d orbitals. Such a reaction will preferentially occur on the exposed surfaces of LSCF where free O2, Sr/La, and M•Mare simultaneously available. As one product of the reaction, SrO(s) is precipitated out on the surfaces of LSCF. The overall process can be described by the following defect reaction (M = Fe and Co) 1 O2 (g) + M•M + Sr /La = SrO(s) + M×M 2

(2)

Since the reaction 1 that produces electron holes M•Mis energetically more favorable, it is reasonable to assume that reaction 2 is the rate-limiting step for the overall surface Srenrichment process. The proposed defect model 2 predicts that higher temperature and pO2 are the favorable conditions to promote the surface Sr-enrichment, which is consistent with all experimental results reported.1−5 On the other hand, electrons from the antibonding σ*-band via the external circuit (under loading) would inhibit reaction 2 by injecting free electrons into the M3d band, reducing hole concentration and thus alleviating the Sr segregation. The behavior that higher cathodic polarization reduces the Sr segregation has indeed been experimentally observed.14 As has been previously pointed out,48 the ORR activity of LSCF cathode is largely determined by the intermediate spin (IS) Co3+ (orCo×Co) on the surface of LSCF. The occupied (3z2−r2) orbital on Co3+ would be oriented toward a surface oxygen vacancy that also attracts an absorbed O2 molecule, followed by an easy transfer of electrons from the single σbonding d orbital to the adsorbed O2 molecule. The formation of Co3+ (orCo×Co) through reaction 2 during the Sr-enrichment process would suggest an enhancement of the ORR activity at least during the early stage of the SrO(s) enrichment (low surface coverage), followed by a deactivation as the SrO(s) coverage increases. This prediction is consistent with recent studies on the surface Sr segregation from LSCo,15,49−51 in which the presence of Sr-enriched phases identified by in situ ambient pressure XPS was found to promote the oxygen surface exchange rate at low coverage, but reduce the activity at higher coverage. Suppression of Surface Sr-Enrichment by the ALDZrO2 Overcoats. The real benefit of the ALD-ZrO2 overcoats on the LSCF-GDC cathode is the improved stability over an extended period of high-temperature operation. If the increase in ASRs of the pristine LSCF-GDC cathode is rooted in the surface enrichment of Sr, the ALD-ZrO2 overcoats shall play an active role in preventing that from happening. As aforementioned, cation interdiffusion occurs between the ZrO2-overcoats and LSCF backbones, the consequence of which is a mutual doping action. The Fe/Co doping in the ZrO2 overcoats can lead to an ORR-promoting mixed oxide-ion and electron conduction, while the Zr doping in LSCF can generate point defect such asZr•Fe(Co). The Zr•Fe(Co) competes with electron holes (M•M) for charge-compensatingSr/La, thus reducing the concentration of electron holes at a fixed doping concentration

(1)

2M•M

where and represent, respectively, two electrons added to the mixed-valent M4+/M3+ redox couple by the introduction of an oxygen vacancy V•• O and restoration of two holes to the M4+/M3+ redox couple by absorption of O2 at V•• O; OO× is a regular lattice oxide ion. Pinning of the Fe4+/Fe3+, Co4+/Co3+, or Ni4+/Ni3+ redox couple at the top of the O-2p bands prevents the surface reaction 1 from being biased completely to the right-hand side in air at the operating temperature Top of the fuel cell. As a result, surface oxygen vacancies are retained in the cathode without eliminating all the mobile holes in the M4+/M3+ couple, which makes the oxide a mixed oxide ion/electronic conductor (MIEC) in air at Top. 4229

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[Sr/La]. Because electron holes are the driver for the surface Srenrichment, the reduction of electron−hole concentration induced by Zr-doping suppresses the Sr-enrichment. No compounds between Sr/La and Zr are expected to form during 800 °C operation from a thermodynamic perspective, which has also been verified by EDS analysis in this study.

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CONCLUSIONS In summary, we demonstrate that the long-term stability of a commercial LSCF-GDC IT-cathode can be significantly improved by creating a conformal layer of ultrathin ZrO2 films on the surface of LSCF grains with ALD. Over a >1100-h testing period at 800 °C, the overcoated cathode shows respective polarization and ohmic ASRs 3 and 1.5 times lower than the pristine sample, whereas the pristine LSCFGDC cathode degrades at a rate 4 times faster than the overcoated one. The multifunctionality of porosity, mixed conductivity and suppressed Sr-enrichment enabled by the ALD-ZrO2 overcoats during high-temperature operation are attributed to the retention of ORR activity. With a defectchemistry model, the suppressed Sr-enrichment is interpreted to be a result of the charge-compensation of Sr/La by ZrM•.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +1-803-777-0204. Fax: +1803-777-0106. Author Contributions

All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The work is supported by the U.S. Army Research Laboratory and the U.S. Army Research Office under Grant W911NF-10R-006, W911NF-13-1-0158, and NSF-CBET1264706. J.B.G. thanks the Robert A. Welch Foundation of Houston, TX, for financial support.

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dx.doi.org/10.1021/cm402879r | Chem. Mater. 2013, 25, 4224−4231

Chemistry of Materials

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