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High‐Performance and Durable Fuel Cells using Co/Sr‐Free Fluorite‐Based Mixed Conducting (Pr,Ce)O2‐δ Cathode

High‐Performance and Durable Fuel Cells using Co/Sr‐Free Fluorite‐Based Mixed Conducting... IntroductionSolid oxide fuel cells (SOFCs) have attracted great interest as environmentally friendly power generation devices due to their high efficiency, fuel flexibility, and low emissions.[1] However, their high operating temperatures (>800 °C) result in accelerated performance degradation, limiting the development and deployment of SOFC technology.[2–5] Over the past several years, significant progress has been made in lowering the operating temperature to less than 650 °C by the use of thin‐film solid electrolytes (hundreds of nm to a few µm) with minimal ohmic loss.[6–10] At the same time, the reduced temperatures lead to exponential increases in the thermally activated electrode polarization resistances, mainly dominated by the oxygen reduction reaction (ORR) at the cathode.[2,11] This makes it particularly challenging to achieve cell performances that match those of conventional SOFCs at these reduced temperatures.Accordingly, extensive studies to date have focused on Co‐containing perovskite mixed ionic electronic conducting (MIEC) oxides, including (La,Sr)(Co,Fe)O3‐δ,[12,13] (Ba,Sr)(Co,Fe)O3‐δ,[3,14] (Sm,Sr)CoO3‐δ,[15,16] Sr(Ti,Fe,Co)O3‐δ,[17] Sr(Y,Nb,Co)O3‐δ,[18] and (Pr,Nd)(Ba,Ca)Co2O5+δ[19,20] given their high MIECs and exceptional ORR activity even at reduced temperatures. For example, Chen et al. fabricated La0.6Sr0.4Co0.2Fe0.8O3[21] and PrBa0.8Ca0.2Co2O5+δ[20] cathodes and obtained exceptionally low area‐specific resistance (ASR) values of 0.43 and ≈0.2 Ω cm2 at 600 °C, respectively. Zhang et al. lowered the ASR of SrTi0.3Fe0.7O3 from 0.65 to 0.21 Ω cm2 at 600 °C by substituting 0.15 Co for Fe on the B‐site.[17] Most successful high‐performance cathodes contain costly transition metal ions (e.g., Co) to achieve high electronic conductivity as well as alkaline earth ions (e.g., Sr) that serve as acceptor elements to achieve simultaneously high electron–hole and oxygen vacancy concentrations.[22] However, unfortunately, Co/Sr‐based perovskite cathodes are also characterized by both chemical and mechanical instabilities during SOFC operations. For example, in addition to its high cost, the variable redox state of Co leads to exceptionally high combined thermal and chemical expansion coefficients (TECs) in the range 20–25 × 10−6 K−1,[18,22–25] limiting their practical use due to incompatibility with typical ceria‐ or zirconia‐based solid electrolytes with much lower expansion coefficients of 10–12 × 10−6 K−1.[26,27] On the other hand, the large diameter of Sr cation used to induce high electronic and ionic carrier densities also leads to preferential segregation of Sr to the perovskite surface and even the formation of second phases, leading to increasingly large charge transfer barriers with time.[28–31] Furthermore, these Sr‐rich phases have been found to react with extrinsic surface impurities such as chromium (Cr), sulfur (S), and silicon (Si) to form insulating chromates, sulfates, and silicates, respectively, that serve to poison the oxygen exchange reaction, leading to long term degradation in performance.[32–34] Identifying highly active and stable cathode materials free of Co and Sr elements should serve to obviate challenges associated with thermal and chemical instability and elevated cost.A number of studies have suggested fluorite‐based Pr‐doped CeO2 (PrxCe1‐xO2‐δ, PCO) as a potential alternative cathode material that exhibits attractive MIEC properties given its surprisingly high levels of oxygen non‐stoichiometry even in the air.[35–42] Furthermore, PCO is free of alkaline earth and transition metal ions such as Co and Sr in the perovskite oxides that were identified as problematic from the standpoint of stability and cost. Moreover, being a binary rather than a ternary oxide like the perovskites, it is less susceptible to sluggish phase transformations and cation nonstoichiometry issues making it also less sensitive to processing conditions. The multiple valence states of Pr (Pr3+ and Pr4+) contribute not only to ionic conduction by the formation of high oxygen vacancy concentrations but also to electron conduction via small polaron hopping through the Pr 4f band, enabling the ORR to occur over the entire PCO surfaces as for other MIECs. Also importantly, its effective TEC value is relatively small (e.g., approximately 13 × 10−6 K−1 for Pr0.2Ce0.8O2‐δ at temperatures below 600 °C[43,44]). This combined with the fact that PCO has the same fluorite crystal structure as the most common solid electrolytes such as yttria‐stabilized zirconia (YSZ) and Gd‐doped ceria (GDC) make it highly compatible with these systems from chemical, thermo‐mechanical and phase stability standpoints. Unfortunately, to date, only significantly lower fuel cell performance (0.098 W cm−2 at 650 °C) has been reported for cells containing PCO cathodes, for example, utilizing a composite Pr0.2Ce0.75Gd0.05O2‐δ‐Gd0.2Ce0.8O2‐δ cathode.[45]Here, we showcase a Co/Sr‐free PCO cathode that both promises enhanced stability and demonstrates the level of high‐performance required for intermediate‐temperature SOFCs. To achieve such enhanced performance, we first systemically examined the surface oxygen exchange coefficient (kchem) of PCO as a function of Pr concentration in PCO thin films with the aid of electrical conductivity relaxation (ECR) measurements. Then, thin‐film electrodes with a vertically‐ordered nanocolumnar structure and selected composition were fabricated by pulsed laser deposition (PLD) to achieve exceptionally high active surface areas. The electrode performance and long‐term stability of these films were then evaluated by electrochemical impedance spectroscopy. First, we found that kchem increases with increasing Pr concentration and those with the highest Pr level studied (x = 0.4) achieved record‐low ASR values of 0.02–0.05 Ω cm2 at 600 °C. Superior long‐term stability for 330 h at 550 °C with degradation of less than 0.2% h−1 was also noted. Finally, we fabricated, for the first time, an anode‐supported cell with a PCO cathode and succeeded in achieving outstanding cell outcomes with a peak power density of 0.92 W cm−2 at 600 °C. These findings demonstrate the feasibility of Co/Sr‐free cathode materials that promise both enhanced chemical/mechanical stability, as well as high performance, for use in intermediate‐temperature SOFCs.Physical and Chemical Characterization of PrxCe1‐xO2‐δ Thin FilmsThe surface oxygen exchange kinetics of mixed conducting oxides is a key characteristic to consider when selecting a cathode material for SOFCs. Here, PLD was selected to fabricate PCO thin films with different Pr contents. Films grown by PLD are advantageous given their ability to analyze surface exchange kinetics accurately and independent of bulk diffusion contributions owing to their dense, flat, and clean surfaces that are free of impurities such as silicon.[46] Approximately 1‐µm‐thick PrxCe1‐xO2‐δ thin films with four different compositions (x = 0.05, 0.1, 0.2, and 0.4), denoted as PCO5, PCO10, PCO20, and PCO40, respectively, were deposited onto c‐Al2O3 (0001) single‐crystal substrates via PLD (Figures S1 and S2, Supporting Information). The chemical compositions of the films were analyzed by X‐ray fluorescence (XRF) and inductively coupled plasma mass spectrometry (ICP‐MS), indicating that the Pr concentrations in the as‐grown films were nearly identical to the nominal compositions (Table S1, Supporting Information). In‐plane and out‐of‐plane high‐resolution X‐ray diffraction (HR‐XRD) analyses suggested that the resulting films are epitaxial with (111)‐orientation relative to the substrate, as shown in Figure S3, Supporting Information. Atomic force microscopy (AFM) and X‐ray photoelectron spectroscopy (XPS) analyses further confirmed that the films were flat (i.e., surface roughness < 1 nm) and Si‐free, as indicated in Figure S4 and Table S2, Supporting Information. Given that grain boundaries may affect the oxygen exchange kinetics in oxygen‐ion conducting oxide, these observations allowed us to ignore potential effects related to grain boundaries and Si surface impurities on the oxygen exchange kinetics.Surface Oxygen Exchange Kinetics of PrxCe1‐xO2‐δ Thin FilmsFigure 1a shows a schematic of the sample arrangement used in the ECR measurements to obtain kchem. Here a PCO thin film and a Pt current collector are sequentially deposited onto an Al2O3 substrate. Scanning electron microscopy (SEM) images show both PCO and Pt surfaces without major pores/voids or cracks, even following ECR measurements. The conductivity relaxation profiles of the PCO films (Figure 1b) show the transient behavior of the conductivity following a rapid change in oxygen partial pressure (pO2) in the surrounding gas. The physical oxygen diffusion length through the films (≈1 µm based on film thickness) is much thinner than the critical thickness (Lc) of PCO,[37] demonstrating that surface oxygen exchange‐controlled kinetics, rather than diffusion‐controlled kinetics, dominate. Accordingly, kchem values were calculated with the aid of Equation (1), as follows:[29,47]1σ(t)−σ(0)σ(∞)−σ(0)=1−exp(−kchemat)\[\begin{array}{*{20}{c}}{\frac{{\sigma \left( t \right) - \sigma \left( 0 \right)}}{{\sigma \left( \infty \right) - \sigma \left( 0 \right)}} = 1 - {\rm{exp}}\left( {\frac{{ - {k_{{\rm{chem}}}}}}{a}t} \right)}\end{array}\]where σ(t) and a represent the electrical conductivity at time t and the thickness of the PCO film, respectively. First, the transient conductivity profiles of PCO films in both reduction and oxidation directions show nearly identical fitting results, reflecting the first‐order surface reaction kinetics (Figure S5, Supporting Information). The oxygen surface exchange coefficient was found to increase monotonically with increasing Pr dopant level. For example, the kchem value of PCO40 was found to be six‐fold higher than that of PCO5 films. Arrhenius plots of kchem, along with the activation energies (Ea) of the PCO films, are shown in Figure 1c. The measured values of kchem and Ea initially depend strongly on the Pr level up to 10%, but then tend to become more weakly dependent on the Pr level above 10% (Figure 1c and Figure S6, Supporting Information). Broadly scattered values for the oxygen surface exchange coefficient of PCO and its activation energies for a given Pr level have been reported in the literature.[35,39] We suspect that these large discrepancies are likely related to the sensitivity that PCO exhibits to preparation methods and small quantities of contaminants such as Si on the surface of PCO.[35,46] In this regard, it is essential to compare the kchem of PCO as a function of Pr content in a single study where other potential contributing factors would not contribute to confusing the trend of kchem on the Pr level. We took great pains to ensure that our prepared PCO films were essentially free of Si and we could thus attribute the changes in kchem and Ea found between the different specimens in this study to come solely from the changes in Pr concentration in our PCO specimens (Table S2, Supporting Information). It is noteworthy that, except for PCO5, our PCO films exhibit acceptable kchem values comparable to those reported for the state‐of‐the‐art perovskite oxides like LSC64[12] as reference. Furthermore, the lower Ea value (≈1.2 eV), compared to LSC64 (1.59 eV), further justifies the use of PCO as an alternative cathode for high‐performance intermediate‐temperature SOFCs.1FigureSurface oxygen exchange kinetics of PCO thin films. a) Illustration of PCO thin films grown on a c‐Al2O3 single‐crystal substrate with Pt thin film current collectors. SEM images of Pt and PCO surfaces after ECR measurements. b) Normalized conductivity transients measured at 600 °C upon switching pO2 from 0.21 to 1 atm. c) Arrhenius plot of kchem with tabulated Ea values for the PCO thin films as a function of Pr concentration. The brown dashed line indicates the state‐of‐the‐art mixed conducting perovskite oxide, La0.6Sr0.4CoO3‐δ (LSC64), as a reference for comparison.[12] d) Comparison of Ea values of kchem with the Pr ionization enthalpy (HPr).[53]Several plausible mechanisms for oxygen exchange on PCO have recently been suggested, but remain under debate in the literature.[35–37,39] For example, Chen et al. proposed that the dissociation of neutral molecular oxygen adsorbate is most likely the rate‐determining step (RDS) at the surface of PCO10 thin films.[36] Schaube et al. found that molecular oxygen species are involved in RDS.[39] They suggested the direct involvement of the redox couple Pr3+/Pr4+ (i.e., related to the formation of the Pr 4f band), likely promoting electron transfer to adsorbed oxygen molecule species. Nicollet et al. instead emphasized the importance of inadvertent surface oxide additives, with different work functions, that can strongly impact the relative surface electron density, and thereby the surface exchange coefficient by modifying the charge transfer reaction.[35] Taken together, the preceding studies suggest that the electronic band structure of PCO is likely a key factor in influencing surface exchange kinetics. As a consequence, we attempted to further examine the electronic band structure of PCO thin films by a combination of UV–vis spectroscopy and UV photoemission spectroscopy (UPS) to investigate how it may have an impact on kchem.[48] We found that the measured energy difference between the Fermi level (EF) and conduction band (ECB) edge decreases as the Pr concentration increases, as shown in Figure S7 and Table S3, Supporting Information. It is important to keep in mind that the Pr impurity band that lies between the valence band with O 2p character and the conduction band with Ce 4f and 5d character, must be quite narrow given the localized character of the Pr 4f energy levels combined with the fact that Pr is diluted in concentration within the Ce cation sublattice. Furthermore, given that the Pr band is partially occupied even in the air (i.e., as evidenced by the formation of some Pr3+), implies that the EF, remains pinned within this rather narrow Pr 4f band. Bishop et al. associated the position of the Pr 4f band relative to the bottom of the conduction band with the enthalpy (HPr) associated with the de‐ionization of an electron from the conduction band back down to the Pr impurity level (PrCe×+e′↔PrCe′\[{\rm{Pr}}_{Ce}^ \times + e' \leftrightarrow {\rm{Pr}}_{Ce}^\prime \]).[42] The narrow character of the Pr 4f band, and the fact that EF is pinned within this band, points to ECB‐EF values being very close to the experimentally determined HPr values previously reported in the literature.[42,49,50] Interestingly, both the magnitude and the dependence of Ea on Pr concentration obtained from kchem values measured by means of conductivity relaxation, appear to be similar in magnitude and Pr dependence to those reported for HPr as shown in Figure 1d. More specifically, as the Pr fraction increases in PCO, HPr decreases, suggesting in turn that EF moves closer to the ECB edge. It is therefore tempting to suggest that a key factor determining the oxygen exchange kinetics on the PCO surface is related to the ease with which electrons are excited to the conduction band, thereby facilitating electron transfer from the conduction band to oxygen molecular adsorbate. A similar correlation between the relative position of EF relative to the ECB edge and the energy (Ea) associated with the oxygen exchange reaction was previously reported by the authors in the MIEC system Sr(Ti1‐xFex)O3.[51] This interpretation is also consistent with the findings of Nicollet et al. that found a strong correlation between the relative surface electron density in PCO, and the kchem that could be varied by orders of magnitude by apparently modulating the band bending at the PCO surface and thus the charge transfer reaction rate.[35]Electrode Design and Electrochemical EvaluationMotivated by the fast oxygen exchange coefficient of PCO comparable to that of LSC64, highly porous, vertically‐ordered, columnar PCO nanostructures were grown onto a single crystal (100) YSZ substrate (8 mol%) under high pressure (100 mTorr O2), offering high specific surface areas (Figure 2a). In‐plane and out‐of‐plane XRD revealed the epitaxial and highly textured nature of the PCO films, in spite of their porous microstructures (Figure S8, Supporting Information). These vertically aligned features (Figure 2e,f) are expected to be advantageous in minimizing potential ohmic losses given the reduced PCO thickness, essential considering the relatively low electronic conductivity of PCO (e.g., σe = 2.8 × 10−2 S cm−1 at 600 °C in PCO20, Figure S9, Supporting Information) compared to that of LSC64 (σe = 2.1 × 103 S cm−1 at 600 °C[12]). Figure 2b–d presents representative SEM images of columnar PCO20 films with respect to the film thickness ranging from ≈0.5 to 2.0 microns. An increase in column height results in a high ratio between the specific and geometric surface area, as indicated in Figure S10, Supporting Information, predicting low ASRs with thicker films.2FigureMicrostructures of vertically‐ordered columnar PCO films. a) Schematic illustration of vertically‐ordered columnar PCO structures and pathway of the oxygen reduction reaction. b–d) Representative SEM images of vertically‐ordered columnar PCO20 films grown onto single‐crystal (100) YSZ substrates at 600 °C under 100 mTorr O2 for a variety of film thicknesses: b) 0.5 µm, c) 1.0 µm, and d) 2.0 µm. e) Cross‐sectional images of columnar films. f) Expanded view of (e) showing the individual columns.The electrode performance capabilities of symmetric cells with the configuration (CC)PCOYSZPCO(CC), where CC designates current collector, were examined by AC impedance spectroscopy between 400–650 °C. Here, two different types of CCs, that is, Ag paste and sputtered metals (Au and Pt), were applied to the vertically‐ordered columnar PCO electrodes. Ag paste was initially used in order to examine the effect of film thickness on ASR. Figure 3a presents a typical impedance spectrum, plotted in Nyquist form, consisting of the offset resistance and a serial semicircle. While the offset resistance is solely attributed to the ohmic resistance of the YSZ electrolyte, the low‐frequency impedance arc reflects the characteristics of the electrochemical reaction occurring at the PCO surface. Enhanced electrode performance was achieved in the 2‐µm‐thick PCO20 sample, which showed a low ASR of ≈0.2 Ω cm2 at 600 °C, in agreement with the previous estimation of the surface area enhancements. It is noteworthy that the intrinsic chemical stability of PCO could bring electrode performance comparable to, or higher than that, of existing perovskite‐based electrodes. The corresponding Ea values (1.2–1.4 eV) compare well to the reported value of 1.26 eV obtained for dense PCO thin films.[37] Furthermore, the calculated capacitances (C) corresponding to the low‐frequency resistances exhibit relatively large values ranging from 37.6 to 238.6 mF cm−2 with −1/6 slope dependence of log C on log pO2 and linear dependence on film thickness (Figure S11, Supporting Information), in line with those previously reported for the volumetric chemical capacitance of mixed conducting PCO.[36–38,41]3FigureElectrochemical performance and stability of columnar PCO electrodes. a) Typical impedance spectra of PCO symmetric cells with Ag paste (CC) grown at 600 °C under 100 mTorr O2 with different film thicknesses (0.5, 1, and 2 µm) obtained at 600 °C, pO2 = 0.21 atm. b) Arrhenius plot of ASRs with regard to thickness, measured with pO2 = 0.21 atm. The gray circle indicates the dense PCO thin films as a reference for comparison.[37] c) Impedance spectra of PCO symmetric cells with Pt sputtering (CC) varying with PLD deposition oxygen pressure (100 and 250 mTorr) and Pr composition (20% Pr and 40% Pr), obtained at 600 °C, pO2 = 0.21 atm. d) Comparison of the electrode performance of PCO cells with those of previously reported Co/Sr‐containing perovskite‐based electrodes fabricated by both PLD[54–56] and typical screen‐printing[16,17,20,57–59] methods. e) Long‐term stability of 2.5‐µm‐thick PCO40 and 1‐µm‐thick LSC64 columnar electrodes measured at 550 °C under dry and wet (2% H2O) air conditions.The effect of working pressure during PLD deposition and Pr composition on the electrode performance was further examined with sputtered Pt CCs. It was found that there is no significant performance difference between Ag paste and sputtered Au and Pt layers, reflecting the fact that the well‐known catalytic properties of Pt did not contribute to the performance outcomes, as indicated in Figure S12, Supporting Information. Based on the literature, higher working pressures typically lead to more random and disordered microstructures, potentially enhancing the specific surface area.[53] As expected, the PCO20 electrode at 250 mTorr O2 outperformed an electrode prepared at lower pressure by a factor of ≈4, as shown in Figure 3c. Given the slightly thicker PCO40 film compared to PCO20, these performance outcomes appear to be similar. Strikingly, for example, the PCO20 film conveys a record‐low ASR of ≈0.05  Ω cm2 at 600 °C, which is highly competitive with those of other benchmark SOFC cathodes, as presented in Figure 3d. This finding shows that PCO films hold great potential as a promising electrode material for high‐performance intermediate‐temperature SOFCs.We subsequently evaluated the stability of the PCO electrode by comparing it to a columnar LSC64 electrode as a reference at 550 °C under both dry and wet (2% H2O) air atmospheres for a prolonged period, as indicated in Figure 3e. In this case, a 1‐µm‐thick columnar LSC64 film layer was deposited at 700 °C under 300 mTorr O2 (Figure S13, Supporting Information). The PCO film displays exceptional stability without notable degradation under both conditions, indicating a low degradation rate of less than 0.2% h−1 for 330 h, as compared to 3.1% h−1 for LSC64 for 200 h (most likely due to Sr segregation, as anticipated). Note that the observed degradation results mainly from the chemical degradation of the electrodes, as confirmed by comparison with offset resistance, including the series resistance, of the symmetric cells (Figures S14 and S15, Supporting Information). These findings confirm our expectation that chemically stable PCO characterized by high ORR activity exhibits considerable potential as an alternative cathode material capable of addressing the long‐standing concerns associated with mixed conducting Co/Sr‐containing perovskite‐based oxides.Demonstration of Anode‐Supported Single CellTo demonstrate the feasibility of PCO as a highly active cathode for intermediate‐temperature SOFCs, a single Ni‐YSZ anode‐supported cell with a YSZ (2 µm) electrolyte and a PCO20 (2 µm) cathode was used to assess its electrochemical performance (Figure 4a). Good contact between the surface of the columnar PCO20 electrode and the Pt thin‐film CC ensures that the entire column participates in the ORR (Figure 4b). The peak power densities of the resulting single cell reach 0.92, 0.60, and 0.28 W cm−2 at 600, 550, and 500 °C, respectively (Figure 4c). To the best of our knowledge, this demonstrates, for the first time, the superior peak power density capabilities of an anode‐supported single cell with a Co/Sr‐free fluorite PCO cathode. Interestingly, it should be noted that the power output even surpasses the record of 0.82, 0.45, and 0.17 W cm−2 achieved at 600, 550, and 500 °C, respectively, by the same platform of an anode‐supported single cell with a state‐of‐the‐art LSC64 cathode as a reference, as indicated in Figure S16, Supporting Information. Furthermore, given the low ORR activation energy associated with the PCO‐based single cell, its performance at reduced operating temperatures improves over that of LSC64 with decreasing temperatures, for example at 500 °C, a remarkable 65% performance enhancement. Figure 4d,e shows that the overall cell performance is determined by the polarization loss of the electrodes and not the ohmic loss of the electrolyte which accounts for only 3% of the total resistance. The ohmic resistance and its Ea value of ≈1.0 eV are consistent with and can therefore be solely attributed to the ionic conductivity of the YSZ electrolyte, even for the case where electrodes like PCO with relatively low electronic conductivity are utilized, as long as such potential ohmic losses can be minimized by the introduction of thin film‐based electrodes. It should be noted that even higher device performance could be achieved by replacing YSZ with more highly conducting ceria‐based electrolytes (e.g., GDC) and the traditional Ni‐YSZ cermet anode with more highly performing anodes (Ni‐GDC). Furthermore, we found no severe degradation of our cell operated under a constant current density of 0.3 A cm−2 at 500 °C. Figure 4f demonstrates the short‐term stability of the operating voltage and power density, exhibiting a degradation rate of only 0.03% h−1 for 10 h. These results point to PCO as being a promising alternative SOFC cathode in addressing the long‐standing chemical stability issues associated with Co/Sr‐based cathodes.4FigureElectrochemical performance and stability of an anode‐supported single cell with columnar PCO cathodes. a) Cross‐sectional SEM image of a Ni‐YSZ anode‐supported single cell with a 2‐µm‐thick YSZ electrolyte layer and a 2‐µm‐thick columnar PCO20 cathode. b) Bright‐field TEM image of the contact between the PCO20 film and Pt CC. c) Typical I–V and I–P curves and d) impedance spectra of the single cell measured at T = 500–600 °C in wet H2 at the anode and air at the cathode. Note that I denotes the current density, V is the cell voltage, and P is the power density. e) Temperature dependence of ASR values corresponding to the total, polarization and ohmic resistance, respectively. f) Short‐term stability test of the single cell for 10 h measured at 500 °C under a constant current density of 0.3 A cm−2.When seeking to demonstrate robust high‐performance SOFCs at intermediate temperatures, several important cathode prerequisites need to be considered: 1) need for adequate levels of MIEC properties to ensure access to the entire oxide surface for the oxygen exchange process to proceed, 2) offer excellent catalytic activities towards ORR, and 3) exhibit high thermal/chemical stability. While the Co/Sr‐based perovskite oxides suffer from criteria 3, the fluorite PCO material suffers instead from low electronic conductivity (criteria 1) that limits current densities. While there have been many attempts over the years to address the chemical instability of Co/Sr‐based cathodes, for example, by modifying the composition and/or surface treatments,[17,23,60,61] the development and optimization of PCO as an electrode remain in their infancy. For the practical application of fuel cells with PCO, most importantly, the electrode has to be designed in such a way as to minimize the ohmic loss induced by the flow of low‐mobility electrons through the PCO layer. For this, first, it is essential to have sufficient Pr concentrations to support the higher level of electronic conductivity. Second, the electrode microstructure should allow electrons transferred from the current collector to be quickly conducted throughout the PCO surface as needed, for example, vertically‐ordered structures with short diffusion lengths as used in this work. Lastly, the micro‐contact between current collectors and PCO structures where ORR can occur on the entire PCO surface is required for the realization of large‐scale SOFCs for high power output. This suggests that there still remains scope to optimize the viable cell design technically for high‐performance and durable fuel cells.Discussion and ConclusionIn this work, we successfully demonstrate the feasibility of PCO as a robust high‐performance cathode for fuel cells by utilizing nanostructured thin‐film‐based PCO electrodes. Given the electrical conductivity of PCO as compared to LSC64 (Table S4, Supporting Information), further optimization, however, is still required to successfully integrate PCO into practical cells for widespread application. Hayd et al., for example, systemically investigated the influence of the microstructure of nano‐scaled LSC thin film cathodes on electrochemical performance by varying processing parameters such as processing temperature, heating rate, and annealing time.[62] As a result, a substantial increase in active surface area was achieved by control of grain size at the nanometric scale (≈17 nm) coupled with high porosity (45%), successfully leading to exceptionally low polarization resistance (0.023 Ω cm2 at 600 °C). Furthermore, the authors pointed out that the choice of the microstructure of the current collector could be optimized to minimize ohmic losses associated with low lateral electronic transport, an issue of concern in PCO. A popular alternative approach for optimizing current collection is the utilization of composite cathode structures. For example, in conventional SOFCs, (La, Sr)MnO3‐δ(LSM)‐YSZ composite structures are extensively used as cathodes given the high electronic and ionic conductivities of LSM and YSZ, respectively. By replacing YSZ, a pure ionic conductor, with MIEC PCO, it can be expected that, for example, a PCO/LSM composite would be much more active given that the ORR would not be limited to only the triple phase boundaries between LSM and YSZ but across the full PCO surface. With this in mind, further microstructural optimization of PCO can be expected in the future to achieve acceptable electrode performance in more conventional electrode configurations.From a scale‐up point of view, the PLD fabrication technique is not a representative method for scaling up this columnar electrode layer. However, columnar structures can be also fabricated via two representative methods for scaling up the electrode layers, which are magnetron sputtering[63] and chemical vapor deposition (CVD).[64] For example, Lee et al. successfully demonstrated an exceptional peak power density of 2.5 W cm−2 at 650 °C by employing columnar structures of the La0.6Sr0.4Co0.2Fe0.8O2.95‐YSZ cathode by a sputtering process.[63] In addition, Oh et al. also fabricated porous columnar thin films of undoped ceria with nano‐scaled grain size by metal‐organic CVD.[64] With these findings, suggested columnar structures, in themselves, should not restrict scale‐up, rather providing others with a guideline to further optimize fabrication processes.Furthermore, from the standpoint of cost and availability, cobalt (Co) and its strategic role presently being played in lithium battery technology today, driven by very rapid growth in electric vehicles, is likely to suffer rapid cost increases and supply disruptions in the not‐too‐distant future. In this regard, the development of Co‐free SOFC cathode material is strongly warranted for next‐generation fuel cells. By comparison, the price of the majority species in PCO, cerium oxide is approximately 25 times lower in cost than that of Co.[65] On the other hand, praseodymium oxide is presently comparable in cost to Co.[66] Based on our findings, while 40% Pr substitution for Ce gives the optimum performance, performance saturates above 10% substitution, offering substantial potential cost savings with little loss in performance.In summary, chemically stable and Co/Sr‐free fluorite‐based mixed conducting PCO was demonstrated, for the first time, to show exceptional intermediate temperature fuel cell performance, importantly coupled with much extended long‐term electrode durability. Electrical conductivity transient profiles confirmed that the kchem increases with Pr concentration, exhibiting a value of 3 × 10−5 cm s−1 for PCO20, comparable to the state‐of‐the‐art perovskite LSC64. The Pr ionization energy is proposed to serve as a key factor impacting the surface exchange kinetics, by impacting electron transfer from the PCO surface to the oxygen molecular adsorbates. Further, the vertically‐ordered, high surface area, columnar PCO structures, not only deliver outstanding electrode activity with an ASR lower than 0.1 Ω cm2 at 600 °C, but also a 15‐fold lower degradation rate of ≈0.2% h‐1 for 330 h compared to that of LSC64. Imposing peak power densities of 0.92 and 0.60 W cm−2 at 600 and 550 °C, respectively, were also achieved using a Ni‐YSZ anode‐supported single cell. Investigating and optimizing PCO as a SOFC cathode, therefore, serves as a promising vehicle for developing durable and high‐performance SOFCs that operate at intermediate temperatures.Experimental SectionFabrication of Dense and Vertically‐Ordered Columnar PCO Thin FilmsDense epitaxial PCO thin films with a thickness of ≈1 µm were grown on single‐crystal Al2O3 (0001) substrates (10 × 10 × 0.5 mm3, MTI Corporation) by PLD, operated with a KrF 248 nm excimer laser emitting at 248 nm (Coherent COMPex Pro 205) at an energy level of 300 mJ with a repetition rate of 10 Hz. PCO targets with different Pr concentrations were prepared by a combined EDTA‐citrate complexing method.[38] During the deposition step, the temperature and the working pressure were 650 °C and 10 mTorr O2, respectively. After the deposition process, the films were annealed at the same temperature under 1 Torr O2 for 20 min to ensure more complete oxidation of the films.Vertically‐ordered columnar PCO thin films were grown on both sides of single‐crystal YSZ (001) substrates (8 mol%, 10 × 10 × 0.5 mm3, MTI Corporation) by PLD (300 mJ, 10 Hz) with the same targets as indicated above. The deposition temperature was 600 °C and the working pressure was 100 and 250 mTorr O2. The film thickness varied from 0.5 to 2.5 µm. With a columnar LSC64 film as a reference for a comparison of the performance and stability, a dense Gd0.1Ce0.9O1.95 (GDC) buffer layer was initially deposited on both sides of the YSZ substrate by PLD at 700 °C under 10 mTorr O2 to prevent any unwanted reaction of YSZ with LSC64. The 1‐µm thick columnar LSC64 film was subsequently deposited on both sides of the GDC/YSZ substrate at 700 °C under 300 mTorr O2.Physical and Chemical CharacterizationThe microstructures of the deposited films were characterized using SEM (Hitachi S‐4800) and cross‐sectional scanning transmission electron microscopy (STEM, Titan cubed G2 60–300, FEI Co.). HR‐XRD (X'pert‐PRO MRD, PANalytical) measurements were taken for both in‐plane and out‐of‐plane reflections of the deposited PCO films using Kα (Cu) radiation (45 kV, 40 mA). The chemical composition of the PCO films was analyzed by an ICP‐MS and XRF. AFM with a Bruker (Innova) device in tapping mode and XPS (K‐alpha, Thermo VG Scientific) were used to investigate the surface roughness and the Si impurities of the films, respectively. The UV–vis spectrum was examined using a UV–vis spectrophotometer (Cary‐300, VARIAN) and UPS with a source energy of 21.21 eV was used to measure the energy difference between the conduction band and the Fermi level of the films.Electrical Conductivity Relaxation MeasurementsECR measurement was conducted by using two‐probe electrodes. With the aid of thin films used in this work, their geometric factor makes the lateral resistance through the films significantly large so that the contact resistance at electrodes should be negligible. This allows the measured resistance to be strongly dependent on the film, not at electrodes (e.g., contact resistance). Two Pt electrodes (200 nm thick and 1 mm width) as current collectors were prepared by DC magnetron sputtering (DC power of 100 W and Ar working pressure of 5 mTorr) with the aid of a shadow mask. The measurements were conducted in an alumina tube at temperatures of 600–650 °C with pO2 steps of 0.21 to 1 atm delivered using a four‐way valve via mass‐flow controllers (MFCs, Fujikin). In‐plane conductivity, which reflects the oxygen content in the PCO thin films, was monitored by measuring the voltage for every 0.5 s to acquire a reliable signal by means of chronopotentiometry (CP, VSP‐300, Biologic) until the sample adapted to a new equilibrium state. The normalized conductivity as a function of time was fitted using the first‐order Equation (1) to calculate the kchem.Electrochemical Measurements of PCO Symmetric CellsThe electrochemical performance of PCO symmetric cells was investigated by AC impedance spectroscopy (ACIS, VSP‐300, Biologic). Two types of current collectors were used: Sputtered metals (Pt and Au) and Ag paste. The cells were placed inside of a continuous‐flow alumina tube for the ACIS test under gas mixtures of O2 and Ar flowed through digital MFCs. Impedance spectra were obtained at temperatures ranging from 400 to 650 °C during the pO2 steps between 0.05 and 1 atm with AC perturbation of 20 mV in a frequency range of 2 MHz to 4 mHz. The performance stability of PCO and LSC64 symmetric cells was assessed for approximately 330 and 200 h, respectively, at 550 °C in dry and wet air atmospheres.Anode‐Supported Single Cell Fabrication and Electrochemical CharacterizationA vertically‐ordered columnar PCO20 cathode (2‐µm thick) was deposited by PLD on a commercial anode‐supported fuel cell with a thin YSZ (2‐µm thick) electrolyte (Elcogen Co.HC400B). The detailed deposition conditions for the PCO cathode are described above. A sputtered Pt current collector was then applied to the PCO cathode. A 2‐µm thick LSC64 cathode layer as a reference for comparison with a PCO‐based single cell was deposited on the GDC/YSZ substrate at 700 °C under 200 mTorr O2. The effective area of the cathode was 1 × 1 cm2. The cell testing configuration for the optimum current collection consisted of a metallic interconnect with a modified rib design and Ni‐foams and Au meshes.[67] Humidified hydrogen (3% H2O–97% H2) and air were applied as the fuel and oxidant to the anode and cathode, respectively. Impedance spectra were measured under open‐circuit voltage in a frequency range of 1 MHz to 0.1 Hz with an AC amplitude of 50 mV. The typical I–V and I–P curves of the single cells were recorded at operating temperatures varying from 600 to 500 °C at intervals of 50 °C using an Iviumstat electrochemical analyzer (Iviumstat, Ivium Technologies). The stability of the voltage was monitored under a constant current density of 0.3 A cm−2 at 500 °C for 10 h.AcknowledgementsThis work was supported by a National Research Foundation of Korea (NRF) grant funded by the Korean government (MSIT) (2021M3H4A1A01002695). J.‐W.S. and D.H.K. also appreciate financial support from Korea Institute of Energy Technology Evaluation and Planning (KETEP), the Ministry of Trade, Industry & Energy, Republic of Korea (No. 20213030030040). H.L.T. acknowledges support from U.S. Department of Energy (DOE), National Energy Technology Laboratory (NETL), Office of Fossil Energy under Award no. DE‐FE0031668.Conflict of InterestThe authors declare no conflict of interest.Author ContributionsH.G.S. and W.J. conceived the idea and designed the experimental protocol. H.G.S. performed the overall sample preparations, characterizations, electrical conductivity relaxation, and electrochemical measurements. D.H.K. and J.‐W.S. performed the fuel cell tests and helped to interpret the results. J.S. helped in collecting the conductivity relaxation data and assisted in the interpretation of the results. S.J.J. assisted with the collection of the thin film characterization data. J.K. assisted in the sample preparation. W.J. supervised the work and provided guidance throughout the project. H.L.T. assisted with the interpretation of the results. H.G.S., J.‐W.S., and W.J. wrote the manuscript with inputs from all co‐authors. All co‐authors contributed by discussing the results and all helped to revise the manuscript.Data Availability StatementResearch data are not shared.B. C. H. Steele, A. Heinzel, 2001, 414, 345.N. P. Brandon, S. Skinner, B. C. H. Steele, Annu. Rev. Mater. Res. 2003, 33, 183.Z. Shao, S. M. Halle, Nature 2004, 431, 170.H.‐I. Ji, J.‐H. Lee, J.‐W. Son, K. J. Yoon, S. Yang, B.‐K. Kim, J. Korean Ceram. Soc. 2020, 57, 480.S. Im, J.‐H. Lee, H.‐I. Ji, J. Korean Ceram. Soc. 2021, 58, 351.A. Evans, A. Bieberle‐Hütter, J. L. M. Rupp, L. J. Gauckler, J. Power Sources 2009, 194, 119.D. Beckel, A. Bieberle‐Hütter, A. Harvey, A. Infortuna, U. P. Muecke, M. Prestat, J. L. M. Rupp, L. J. Gauckler, J. Power Sources 2007, 173, 325.J. An, J. H. Shim, Y.‐B. Kim, J. S. Park, W. Lee, T. M. Gür, F. B. Prinz, MRS Bull. 2014, 39, 798.C.‐W. Kwon, J.‐W. Son, J.‐H. Lee, H.‐M. Kim, H.‐W. Lee, K.‐B. Kim, Adv. Funct. Mater. 2011, 21, 1154.M. Tsuchiya, B. K. Lai, S. Ramanathan, Nat. Nanotechnol. 2011, 6, 282.S. B. Adler, Chem. Rev. 2004, 104, 4791.A. Egger, E. Bucher, M. Yang, W. Sitte, Solid State Ionics 2012, 225, 55.Y. Chen, Y. Choi, S. Yoo, Y. Ding, R. Yan, K. Pei, C. Qu, L. Zhang, I. Chang, B. Zhao, Y. Zhang, H. Chen, Y. Chen, C. Yang, B. deGlee, R. Murphy, J. Liu, M. Liu, Joule 2018, 2, 938.E. Bucher, A. Egger, P. Ried, W. Sitte, P. Holtappels, Solid State Ionics 2008, 179, 1032.C. Xia, W. Rauch, F. Chen, M. Liu, Solid State Ionics 2002, 149, 11.S. W. Baek, J. H. Kim, J. Bae, Solid State Ionics 2008, 179, 1570.S.‐L. Zhang, H. Wang, M. Y. Lu, A.‐P. Zhang, L. V Mogni, Q. Liu, C.‐X. Li, C.‐J. Li, S. A. Barnett, Energy Environ. Sci. 2018, 11, 1870.Y. Zhang, B. Chen, D. Guan, M. Xu, R. Ran, M. Ni, W. Zhou, R. O'Hayre, Z. Shao, Nature 2021, 591, 246.S. Choi, S. Park, J. Shin, G. Kim, J. Mater. Chem. A 2015, 3, 6088.Y. Chen, S. Yoo, Y. Choi, J. H. Kim, Y. Ding, K. Pei, R. Murphy, Y. Zhang, B. Zhao, W. Zhang, H. Chen, Y. Chen, W. Yuan, C. Yang, M. Liu, Energy Environ. Sci. 2018, 11, 2458.Y. Chen, Y. Bu, Y. Zhang, R. Yan, D. Ding, B. Zhao, S. Yoo, D. Dang, R. Hu, C. Yang, M. Liu, Adv. Energy Mater. 2017, 7, 1601890.F. Prado, T. Armstrong, A. Caneiro, A. Manthiram, J. Electrochem. Soc. 2001, 148, J7.J. Wang, K. Y. Lam, M. Saccoccio, Y. Gao, D. Chen, F. Ciucci, J. Power Sources 2016, 324, 224.S. Hou, J. A. Alonso, J. B. Goodenough, J. Power Sources 2010, 195, 280.F. Wang, Q. Zhou, T. He, G. Li, H. Ding, J. Power Sources 2010, 195, 3772.H. Hayashi, M. Kanoh, C. J. Quan, H. Inaba, S. Wang, M. Dokiya, H. Tagawa, Solid State Ionics 2000, 132, 227.H. Hayashi, T. Saitou, N. Maruyama, H. Inaba, K. Kawamura, M. Mori, Solid State Ionics 2005, 176, 613.B. Koo, K. Kim, J. K. Kim, H. Kwon, J. W. Han, W. C. Jung, Joule 2018, 2, 1476.B. Koo, H. Kwon, Y. Kim, H. G. Seo, J. W. Han, W. Jung, Energy Environ. Sci. 2018, 11, 71.Z. Cai, M. Kubicek, J. Fleig, B. Yildiz, Chem. Mater. 2012, 24, 1116.W. Jung, H. L. Tuller, Energy Environ. Sci. 2012, 5, 5370.C. C. Wang, M. Gholizadeh, B. Hou, X. Fan, RSC Adv. 2021, 11, 7.E. Bucher, W. Sitte, Solid State Ionics 2011, 192, 480.A. F. Staerz, H. G. Seo, T. Defferriere, H. L. Tuller, J. Mater. Chem. A 2022, 10, 2618.C. Nicollet, C. Toparli, G. F. Harrington, T. Defferriere, B. Yildiz, H. L. Tuller, Nat. Catal. 2020, 3, 913.D. Chen, Z. Guan, D. Zhang, L. Trotochaud, E. Crumlin, S. Nemsak, H. Bluhm, H. L. Tuller, W. C. Chueh, Nat. Catal. 2020, 3, 116.D. Chen, S. R. Bishop, H. L. Tuller, J. Electroceram. 2012, 28, 62.H. G. Seo, Y. Choi, W. C. Jung, Adv. Energy Mater. 2018, 8, 1703647.M. Schaube, R. Merkle, J. Maier, J. Mater. Chem. A 2019, 7, 21854.H. Kim, H. G. Seo, Y. Choi, D.‐K. Lim, W. Jung, J. Mater. Chem. A 2020, 8, 14491.D. Chen, S. R. Bishop, H. L. Tuller, Adv. Funct. Mater. 2013, 23, 2168.S. R. Bishop, T. S. Stefanik, H. L. Tuller, Phys. Chem. Chem. Phys. 2011, 13, 10165.D. P. Fagg, V. V Kharton, A. Shaula, I. P. Marozau, J. R. Frade, Solid State Ionics 2005, 176, 1723.C. Lenser, F. Gunkel, Y. J. Sohn, N. H. Menzler, Solid State Ionics 2018, 314, 204.R. Chockalingam, A. K. Ganguli, S. Basu, J. Power Sources 2014, 250, 80.L. Zhao, N. H. Perry, T. Daio, K. Sasaki, S. R. Bishop, Chem. Mater. 2015, 27, 3065.C. B. Gopal, S. M. Haile, J. Mater. Chem. A 2014, 2, 2405.Y. Yin, S. Fu, S. Zhou, Y. Song, L. Li, M. Zhang, J. Wang, P. Mariyappan, S. M. Alshehri, T. Ahamad, Y. Yamauchi, Electron. Mater. Lett. 2020, 16, 224.J. J. Kim, S. R. Bishop, N. Thompson, Y. Kuru, H. L. Tuller, Solid State Ionics 2012, 225, 198.K. Schmale, M. Grünebaum, M. Janssen, S. Baumann, F. Schulze‐Küppers, H.‐D. Wiemhöfer, Phys. Status Solidi 2011, 248, 314.W. Jung, H. L. Tuller, Adv. Energy Mater. 2011, 1, 1184.T. S. Stefanik, Electrical Properties and Defect Structure of Praseodymium‐Cerium Oxide Solid Solutions, Ph. D., Massachusetts Institute of Technology, Cambridge, MA 2004.A. Infortuna, A. S. Harvey, L. J. Gauckler, Adv. Funct. Mater. 2008, 18, 127.J.‐H. Park, W.‐S. Hong, K. J. Yoon, J.‐H. Lee, H.‐W. Lee, J.‐W. Son, J. Electrochem. Soc. 2014, 161, F16.D. Beckel, U. P. Muecke, T. Gyger, G. Florey, A. Infortuna, L. J. Gauckler, Solid State Ionics 2007, 178, 407.J. Yoon, R. Araujo, N. Grunbaum, L. Baqué, A. Serquis, A. Caneiro, X. Zhang, H. Wang, Appl. Surf. Sci. 2007, 254, 266.M. Shang, J. Tong, R. O'Hayre, RSC Adv. 2013, 3, 15769.C. Duan, D. Hook, Y. Chen, J. Tong, R. O'Hayre, Energy Environ. Sci. 2017, 10, 176.M. Li, M. Zhao, F. Li, W. Zhou, V. K. Peterson, X. Xu, Z. Shao, I. Gentle, Z. Zhu, Nat. Commun. 2017, 8, 1.B. Koo, J. Seo, J. K. Kim, W. Jung, J. Mater. Chem. A 2020, 8, 13763.N. Tsvetkov, Q. Lu, L. Sun, E. J. Crumlin, B. Yildiz, Nat. Mater. 2016, 15, 1010.J. Hayd, L. Dieterle, U. Guntow, D. Gerthsen, E. Ivers‐Tiffée, J. Power Sources 2011, 196, 7263.Y. H. Lee, H. Ren, E. A. Wu, E. E. Fullerton, Y. S. Meng, N. Q. Minh, Nano Lett. 2020, 20, 2943.T.‐S. Oh, D. A. Boyd, D. G. Goodwin, S. M. Haile, Phys. Chem. Chem. Phys. 2013, 15, 2466.U.S. Geological Survey, Mineral Commodity Summaries 2022, U.S. Geological Survey, Reston, VA, 2022.Statista, Praseodymium Oxide Price Worldwide from 2009 to 2020 with a Forecast from 2021 to 2030 (in U.S. Dollars per Metric Ton), Stormcrow, 2021.H. S. Noh, J. Hwang, K. Yoon, B. K. Kim, H. W. Lee, J. H. Lee, J. W. Son, J. Power Sources 2013, 230, 109. http://www.deepdyve.com/assets/images/DeepDyve-Logo-lg.png Advanced Energy Materials Wiley

High‐Performance and Durable Fuel Cells using Co/Sr‐Free Fluorite‐Based Mixed Conducting (Pr,Ce)O2‐δ Cathode

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Wiley
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© 2022 Wiley‐VCH GmbH
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1614-6832
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1614-6840
DOI
10.1002/aenm.202202101
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Abstract

IntroductionSolid oxide fuel cells (SOFCs) have attracted great interest as environmentally friendly power generation devices due to their high efficiency, fuel flexibility, and low emissions.[1] However, their high operating temperatures (>800 °C) result in accelerated performance degradation, limiting the development and deployment of SOFC technology.[2–5] Over the past several years, significant progress has been made in lowering the operating temperature to less than 650 °C by the use of thin‐film solid electrolytes (hundreds of nm to a few µm) with minimal ohmic loss.[6–10] At the same time, the reduced temperatures lead to exponential increases in the thermally activated electrode polarization resistances, mainly dominated by the oxygen reduction reaction (ORR) at the cathode.[2,11] This makes it particularly challenging to achieve cell performances that match those of conventional SOFCs at these reduced temperatures.Accordingly, extensive studies to date have focused on Co‐containing perovskite mixed ionic electronic conducting (MIEC) oxides, including (La,Sr)(Co,Fe)O3‐δ,[12,13] (Ba,Sr)(Co,Fe)O3‐δ,[3,14] (Sm,Sr)CoO3‐δ,[15,16] Sr(Ti,Fe,Co)O3‐δ,[17] Sr(Y,Nb,Co)O3‐δ,[18] and (Pr,Nd)(Ba,Ca)Co2O5+δ[19,20] given their high MIECs and exceptional ORR activity even at reduced temperatures. For example, Chen et al. fabricated La0.6Sr0.4Co0.2Fe0.8O3[21] and PrBa0.8Ca0.2Co2O5+δ[20] cathodes and obtained exceptionally low area‐specific resistance (ASR) values of 0.43 and ≈0.2 Ω cm2 at 600 °C, respectively. Zhang et al. lowered the ASR of SrTi0.3Fe0.7O3 from 0.65 to 0.21 Ω cm2 at 600 °C by substituting 0.15 Co for Fe on the B‐site.[17] Most successful high‐performance cathodes contain costly transition metal ions (e.g., Co) to achieve high electronic conductivity as well as alkaline earth ions (e.g., Sr) that serve as acceptor elements to achieve simultaneously high electron–hole and oxygen vacancy concentrations.[22] However, unfortunately, Co/Sr‐based perovskite cathodes are also characterized by both chemical and mechanical instabilities during SOFC operations. For example, in addition to its high cost, the variable redox state of Co leads to exceptionally high combined thermal and chemical expansion coefficients (TECs) in the range 20–25 × 10−6 K−1,[18,22–25] limiting their practical use due to incompatibility with typical ceria‐ or zirconia‐based solid electrolytes with much lower expansion coefficients of 10–12 × 10−6 K−1.[26,27] On the other hand, the large diameter of Sr cation used to induce high electronic and ionic carrier densities also leads to preferential segregation of Sr to the perovskite surface and even the formation of second phases, leading to increasingly large charge transfer barriers with time.[28–31] Furthermore, these Sr‐rich phases have been found to react with extrinsic surface impurities such as chromium (Cr), sulfur (S), and silicon (Si) to form insulating chromates, sulfates, and silicates, respectively, that serve to poison the oxygen exchange reaction, leading to long term degradation in performance.[32–34] Identifying highly active and stable cathode materials free of Co and Sr elements should serve to obviate challenges associated with thermal and chemical instability and elevated cost.A number of studies have suggested fluorite‐based Pr‐doped CeO2 (PrxCe1‐xO2‐δ, PCO) as a potential alternative cathode material that exhibits attractive MIEC properties given its surprisingly high levels of oxygen non‐stoichiometry even in the air.[35–42] Furthermore, PCO is free of alkaline earth and transition metal ions such as Co and Sr in the perovskite oxides that were identified as problematic from the standpoint of stability and cost. Moreover, being a binary rather than a ternary oxide like the perovskites, it is less susceptible to sluggish phase transformations and cation nonstoichiometry issues making it also less sensitive to processing conditions. The multiple valence states of Pr (Pr3+ and Pr4+) contribute not only to ionic conduction by the formation of high oxygen vacancy concentrations but also to electron conduction via small polaron hopping through the Pr 4f band, enabling the ORR to occur over the entire PCO surfaces as for other MIECs. Also importantly, its effective TEC value is relatively small (e.g., approximately 13 × 10−6 K−1 for Pr0.2Ce0.8O2‐δ at temperatures below 600 °C[43,44]). This combined with the fact that PCO has the same fluorite crystal structure as the most common solid electrolytes such as yttria‐stabilized zirconia (YSZ) and Gd‐doped ceria (GDC) make it highly compatible with these systems from chemical, thermo‐mechanical and phase stability standpoints. Unfortunately, to date, only significantly lower fuel cell performance (0.098 W cm−2 at 650 °C) has been reported for cells containing PCO cathodes, for example, utilizing a composite Pr0.2Ce0.75Gd0.05O2‐δ‐Gd0.2Ce0.8O2‐δ cathode.[45]Here, we showcase a Co/Sr‐free PCO cathode that both promises enhanced stability and demonstrates the level of high‐performance required for intermediate‐temperature SOFCs. To achieve such enhanced performance, we first systemically examined the surface oxygen exchange coefficient (kchem) of PCO as a function of Pr concentration in PCO thin films with the aid of electrical conductivity relaxation (ECR) measurements. Then, thin‐film electrodes with a vertically‐ordered nanocolumnar structure and selected composition were fabricated by pulsed laser deposition (PLD) to achieve exceptionally high active surface areas. The electrode performance and long‐term stability of these films were then evaluated by electrochemical impedance spectroscopy. First, we found that kchem increases with increasing Pr concentration and those with the highest Pr level studied (x = 0.4) achieved record‐low ASR values of 0.02–0.05 Ω cm2 at 600 °C. Superior long‐term stability for 330 h at 550 °C with degradation of less than 0.2% h−1 was also noted. Finally, we fabricated, for the first time, an anode‐supported cell with a PCO cathode and succeeded in achieving outstanding cell outcomes with a peak power density of 0.92 W cm−2 at 600 °C. These findings demonstrate the feasibility of Co/Sr‐free cathode materials that promise both enhanced chemical/mechanical stability, as well as high performance, for use in intermediate‐temperature SOFCs.Physical and Chemical Characterization of PrxCe1‐xO2‐δ Thin FilmsThe surface oxygen exchange kinetics of mixed conducting oxides is a key characteristic to consider when selecting a cathode material for SOFCs. Here, PLD was selected to fabricate PCO thin films with different Pr contents. Films grown by PLD are advantageous given their ability to analyze surface exchange kinetics accurately and independent of bulk diffusion contributions owing to their dense, flat, and clean surfaces that are free of impurities such as silicon.[46] Approximately 1‐µm‐thick PrxCe1‐xO2‐δ thin films with four different compositions (x = 0.05, 0.1, 0.2, and 0.4), denoted as PCO5, PCO10, PCO20, and PCO40, respectively, were deposited onto c‐Al2O3 (0001) single‐crystal substrates via PLD (Figures S1 and S2, Supporting Information). The chemical compositions of the films were analyzed by X‐ray fluorescence (XRF) and inductively coupled plasma mass spectrometry (ICP‐MS), indicating that the Pr concentrations in the as‐grown films were nearly identical to the nominal compositions (Table S1, Supporting Information). In‐plane and out‐of‐plane high‐resolution X‐ray diffraction (HR‐XRD) analyses suggested that the resulting films are epitaxial with (111)‐orientation relative to the substrate, as shown in Figure S3, Supporting Information. Atomic force microscopy (AFM) and X‐ray photoelectron spectroscopy (XPS) analyses further confirmed that the films were flat (i.e., surface roughness < 1 nm) and Si‐free, as indicated in Figure S4 and Table S2, Supporting Information. Given that grain boundaries may affect the oxygen exchange kinetics in oxygen‐ion conducting oxide, these observations allowed us to ignore potential effects related to grain boundaries and Si surface impurities on the oxygen exchange kinetics.Surface Oxygen Exchange Kinetics of PrxCe1‐xO2‐δ Thin FilmsFigure 1a shows a schematic of the sample arrangement used in the ECR measurements to obtain kchem. Here a PCO thin film and a Pt current collector are sequentially deposited onto an Al2O3 substrate. Scanning electron microscopy (SEM) images show both PCO and Pt surfaces without major pores/voids or cracks, even following ECR measurements. The conductivity relaxation profiles of the PCO films (Figure 1b) show the transient behavior of the conductivity following a rapid change in oxygen partial pressure (pO2) in the surrounding gas. The physical oxygen diffusion length through the films (≈1 µm based on film thickness) is much thinner than the critical thickness (Lc) of PCO,[37] demonstrating that surface oxygen exchange‐controlled kinetics, rather than diffusion‐controlled kinetics, dominate. Accordingly, kchem values were calculated with the aid of Equation (1), as follows:[29,47]1σ(t)−σ(0)σ(∞)−σ(0)=1−exp(−kchemat)\[\begin{array}{*{20}{c}}{\frac{{\sigma \left( t \right) - \sigma \left( 0 \right)}}{{\sigma \left( \infty \right) - \sigma \left( 0 \right)}} = 1 - {\rm{exp}}\left( {\frac{{ - {k_{{\rm{chem}}}}}}{a}t} \right)}\end{array}\]where σ(t) and a represent the electrical conductivity at time t and the thickness of the PCO film, respectively. First, the transient conductivity profiles of PCO films in both reduction and oxidation directions show nearly identical fitting results, reflecting the first‐order surface reaction kinetics (Figure S5, Supporting Information). The oxygen surface exchange coefficient was found to increase monotonically with increasing Pr dopant level. For example, the kchem value of PCO40 was found to be six‐fold higher than that of PCO5 films. Arrhenius plots of kchem, along with the activation energies (Ea) of the PCO films, are shown in Figure 1c. The measured values of kchem and Ea initially depend strongly on the Pr level up to 10%, but then tend to become more weakly dependent on the Pr level above 10% (Figure 1c and Figure S6, Supporting Information). Broadly scattered values for the oxygen surface exchange coefficient of PCO and its activation energies for a given Pr level have been reported in the literature.[35,39] We suspect that these large discrepancies are likely related to the sensitivity that PCO exhibits to preparation methods and small quantities of contaminants such as Si on the surface of PCO.[35,46] In this regard, it is essential to compare the kchem of PCO as a function of Pr content in a single study where other potential contributing factors would not contribute to confusing the trend of kchem on the Pr level. We took great pains to ensure that our prepared PCO films were essentially free of Si and we could thus attribute the changes in kchem and Ea found between the different specimens in this study to come solely from the changes in Pr concentration in our PCO specimens (Table S2, Supporting Information). It is noteworthy that, except for PCO5, our PCO films exhibit acceptable kchem values comparable to those reported for the state‐of‐the‐art perovskite oxides like LSC64[12] as reference. Furthermore, the lower Ea value (≈1.2 eV), compared to LSC64 (1.59 eV), further justifies the use of PCO as an alternative cathode for high‐performance intermediate‐temperature SOFCs.1FigureSurface oxygen exchange kinetics of PCO thin films. a) Illustration of PCO thin films grown on a c‐Al2O3 single‐crystal substrate with Pt thin film current collectors. SEM images of Pt and PCO surfaces after ECR measurements. b) Normalized conductivity transients measured at 600 °C upon switching pO2 from 0.21 to 1 atm. c) Arrhenius plot of kchem with tabulated Ea values for the PCO thin films as a function of Pr concentration. The brown dashed line indicates the state‐of‐the‐art mixed conducting perovskite oxide, La0.6Sr0.4CoO3‐δ (LSC64), as a reference for comparison.[12] d) Comparison of Ea values of kchem with the Pr ionization enthalpy (HPr).[53]Several plausible mechanisms for oxygen exchange on PCO have recently been suggested, but remain under debate in the literature.[35–37,39] For example, Chen et al. proposed that the dissociation of neutral molecular oxygen adsorbate is most likely the rate‐determining step (RDS) at the surface of PCO10 thin films.[36] Schaube et al. found that molecular oxygen species are involved in RDS.[39] They suggested the direct involvement of the redox couple Pr3+/Pr4+ (i.e., related to the formation of the Pr 4f band), likely promoting electron transfer to adsorbed oxygen molecule species. Nicollet et al. instead emphasized the importance of inadvertent surface oxide additives, with different work functions, that can strongly impact the relative surface electron density, and thereby the surface exchange coefficient by modifying the charge transfer reaction.[35] Taken together, the preceding studies suggest that the electronic band structure of PCO is likely a key factor in influencing surface exchange kinetics. As a consequence, we attempted to further examine the electronic band structure of PCO thin films by a combination of UV–vis spectroscopy and UV photoemission spectroscopy (UPS) to investigate how it may have an impact on kchem.[48] We found that the measured energy difference between the Fermi level (EF) and conduction band (ECB) edge decreases as the Pr concentration increases, as shown in Figure S7 and Table S3, Supporting Information. It is important to keep in mind that the Pr impurity band that lies between the valence band with O 2p character and the conduction band with Ce 4f and 5d character, must be quite narrow given the localized character of the Pr 4f energy levels combined with the fact that Pr is diluted in concentration within the Ce cation sublattice. Furthermore, given that the Pr band is partially occupied even in the air (i.e., as evidenced by the formation of some Pr3+), implies that the EF, remains pinned within this rather narrow Pr 4f band. Bishop et al. associated the position of the Pr 4f band relative to the bottom of the conduction band with the enthalpy (HPr) associated with the de‐ionization of an electron from the conduction band back down to the Pr impurity level (PrCe×+e′↔PrCe′\[{\rm{Pr}}_{Ce}^ \times + e' \leftrightarrow {\rm{Pr}}_{Ce}^\prime \]).[42] The narrow character of the Pr 4f band, and the fact that EF is pinned within this band, points to ECB‐EF values being very close to the experimentally determined HPr values previously reported in the literature.[42,49,50] Interestingly, both the magnitude and the dependence of Ea on Pr concentration obtained from kchem values measured by means of conductivity relaxation, appear to be similar in magnitude and Pr dependence to those reported for HPr as shown in Figure 1d. More specifically, as the Pr fraction increases in PCO, HPr decreases, suggesting in turn that EF moves closer to the ECB edge. It is therefore tempting to suggest that a key factor determining the oxygen exchange kinetics on the PCO surface is related to the ease with which electrons are excited to the conduction band, thereby facilitating electron transfer from the conduction band to oxygen molecular adsorbate. A similar correlation between the relative position of EF relative to the ECB edge and the energy (Ea) associated with the oxygen exchange reaction was previously reported by the authors in the MIEC system Sr(Ti1‐xFex)O3.[51] This interpretation is also consistent with the findings of Nicollet et al. that found a strong correlation between the relative surface electron density in PCO, and the kchem that could be varied by orders of magnitude by apparently modulating the band bending at the PCO surface and thus the charge transfer reaction rate.[35]Electrode Design and Electrochemical EvaluationMotivated by the fast oxygen exchange coefficient of PCO comparable to that of LSC64, highly porous, vertically‐ordered, columnar PCO nanostructures were grown onto a single crystal (100) YSZ substrate (8 mol%) under high pressure (100 mTorr O2), offering high specific surface areas (Figure 2a). In‐plane and out‐of‐plane XRD revealed the epitaxial and highly textured nature of the PCO films, in spite of their porous microstructures (Figure S8, Supporting Information). These vertically aligned features (Figure 2e,f) are expected to be advantageous in minimizing potential ohmic losses given the reduced PCO thickness, essential considering the relatively low electronic conductivity of PCO (e.g., σe = 2.8 × 10−2 S cm−1 at 600 °C in PCO20, Figure S9, Supporting Information) compared to that of LSC64 (σe = 2.1 × 103 S cm−1 at 600 °C[12]). Figure 2b–d presents representative SEM images of columnar PCO20 films with respect to the film thickness ranging from ≈0.5 to 2.0 microns. An increase in column height results in a high ratio between the specific and geometric surface area, as indicated in Figure S10, Supporting Information, predicting low ASRs with thicker films.2FigureMicrostructures of vertically‐ordered columnar PCO films. a) Schematic illustration of vertically‐ordered columnar PCO structures and pathway of the oxygen reduction reaction. b–d) Representative SEM images of vertically‐ordered columnar PCO20 films grown onto single‐crystal (100) YSZ substrates at 600 °C under 100 mTorr O2 for a variety of film thicknesses: b) 0.5 µm, c) 1.0 µm, and d) 2.0 µm. e) Cross‐sectional images of columnar films. f) Expanded view of (e) showing the individual columns.The electrode performance capabilities of symmetric cells with the configuration (CC)PCOYSZPCO(CC), where CC designates current collector, were examined by AC impedance spectroscopy between 400–650 °C. Here, two different types of CCs, that is, Ag paste and sputtered metals (Au and Pt), were applied to the vertically‐ordered columnar PCO electrodes. Ag paste was initially used in order to examine the effect of film thickness on ASR. Figure 3a presents a typical impedance spectrum, plotted in Nyquist form, consisting of the offset resistance and a serial semicircle. While the offset resistance is solely attributed to the ohmic resistance of the YSZ electrolyte, the low‐frequency impedance arc reflects the characteristics of the electrochemical reaction occurring at the PCO surface. Enhanced electrode performance was achieved in the 2‐µm‐thick PCO20 sample, which showed a low ASR of ≈0.2 Ω cm2 at 600 °C, in agreement with the previous estimation of the surface area enhancements. It is noteworthy that the intrinsic chemical stability of PCO could bring electrode performance comparable to, or higher than that, of existing perovskite‐based electrodes. The corresponding Ea values (1.2–1.4 eV) compare well to the reported value of 1.26 eV obtained for dense PCO thin films.[37] Furthermore, the calculated capacitances (C) corresponding to the low‐frequency resistances exhibit relatively large values ranging from 37.6 to 238.6 mF cm−2 with −1/6 slope dependence of log C on log pO2 and linear dependence on film thickness (Figure S11, Supporting Information), in line with those previously reported for the volumetric chemical capacitance of mixed conducting PCO.[36–38,41]3FigureElectrochemical performance and stability of columnar PCO electrodes. a) Typical impedance spectra of PCO symmetric cells with Ag paste (CC) grown at 600 °C under 100 mTorr O2 with different film thicknesses (0.5, 1, and 2 µm) obtained at 600 °C, pO2 = 0.21 atm. b) Arrhenius plot of ASRs with regard to thickness, measured with pO2 = 0.21 atm. The gray circle indicates the dense PCO thin films as a reference for comparison.[37] c) Impedance spectra of PCO symmetric cells with Pt sputtering (CC) varying with PLD deposition oxygen pressure (100 and 250 mTorr) and Pr composition (20% Pr and 40% Pr), obtained at 600 °C, pO2 = 0.21 atm. d) Comparison of the electrode performance of PCO cells with those of previously reported Co/Sr‐containing perovskite‐based electrodes fabricated by both PLD[54–56] and typical screen‐printing[16,17,20,57–59] methods. e) Long‐term stability of 2.5‐µm‐thick PCO40 and 1‐µm‐thick LSC64 columnar electrodes measured at 550 °C under dry and wet (2% H2O) air conditions.The effect of working pressure during PLD deposition and Pr composition on the electrode performance was further examined with sputtered Pt CCs. It was found that there is no significant performance difference between Ag paste and sputtered Au and Pt layers, reflecting the fact that the well‐known catalytic properties of Pt did not contribute to the performance outcomes, as indicated in Figure S12, Supporting Information. Based on the literature, higher working pressures typically lead to more random and disordered microstructures, potentially enhancing the specific surface area.[53] As expected, the PCO20 electrode at 250 mTorr O2 outperformed an electrode prepared at lower pressure by a factor of ≈4, as shown in Figure 3c. Given the slightly thicker PCO40 film compared to PCO20, these performance outcomes appear to be similar. Strikingly, for example, the PCO20 film conveys a record‐low ASR of ≈0.05  Ω cm2 at 600 °C, which is highly competitive with those of other benchmark SOFC cathodes, as presented in Figure 3d. This finding shows that PCO films hold great potential as a promising electrode material for high‐performance intermediate‐temperature SOFCs.We subsequently evaluated the stability of the PCO electrode by comparing it to a columnar LSC64 electrode as a reference at 550 °C under both dry and wet (2% H2O) air atmospheres for a prolonged period, as indicated in Figure 3e. In this case, a 1‐µm‐thick columnar LSC64 film layer was deposited at 700 °C under 300 mTorr O2 (Figure S13, Supporting Information). The PCO film displays exceptional stability without notable degradation under both conditions, indicating a low degradation rate of less than 0.2% h−1 for 330 h, as compared to 3.1% h−1 for LSC64 for 200 h (most likely due to Sr segregation, as anticipated). Note that the observed degradation results mainly from the chemical degradation of the electrodes, as confirmed by comparison with offset resistance, including the series resistance, of the symmetric cells (Figures S14 and S15, Supporting Information). These findings confirm our expectation that chemically stable PCO characterized by high ORR activity exhibits considerable potential as an alternative cathode material capable of addressing the long‐standing concerns associated with mixed conducting Co/Sr‐containing perovskite‐based oxides.Demonstration of Anode‐Supported Single CellTo demonstrate the feasibility of PCO as a highly active cathode for intermediate‐temperature SOFCs, a single Ni‐YSZ anode‐supported cell with a YSZ (2 µm) electrolyte and a PCO20 (2 µm) cathode was used to assess its electrochemical performance (Figure 4a). Good contact between the surface of the columnar PCO20 electrode and the Pt thin‐film CC ensures that the entire column participates in the ORR (Figure 4b). The peak power densities of the resulting single cell reach 0.92, 0.60, and 0.28 W cm−2 at 600, 550, and 500 °C, respectively (Figure 4c). To the best of our knowledge, this demonstrates, for the first time, the superior peak power density capabilities of an anode‐supported single cell with a Co/Sr‐free fluorite PCO cathode. Interestingly, it should be noted that the power output even surpasses the record of 0.82, 0.45, and 0.17 W cm−2 achieved at 600, 550, and 500 °C, respectively, by the same platform of an anode‐supported single cell with a state‐of‐the‐art LSC64 cathode as a reference, as indicated in Figure S16, Supporting Information. Furthermore, given the low ORR activation energy associated with the PCO‐based single cell, its performance at reduced operating temperatures improves over that of LSC64 with decreasing temperatures, for example at 500 °C, a remarkable 65% performance enhancement. Figure 4d,e shows that the overall cell performance is determined by the polarization loss of the electrodes and not the ohmic loss of the electrolyte which accounts for only 3% of the total resistance. The ohmic resistance and its Ea value of ≈1.0 eV are consistent with and can therefore be solely attributed to the ionic conductivity of the YSZ electrolyte, even for the case where electrodes like PCO with relatively low electronic conductivity are utilized, as long as such potential ohmic losses can be minimized by the introduction of thin film‐based electrodes. It should be noted that even higher device performance could be achieved by replacing YSZ with more highly conducting ceria‐based electrolytes (e.g., GDC) and the traditional Ni‐YSZ cermet anode with more highly performing anodes (Ni‐GDC). Furthermore, we found no severe degradation of our cell operated under a constant current density of 0.3 A cm−2 at 500 °C. Figure 4f demonstrates the short‐term stability of the operating voltage and power density, exhibiting a degradation rate of only 0.03% h−1 for 10 h. These results point to PCO as being a promising alternative SOFC cathode in addressing the long‐standing chemical stability issues associated with Co/Sr‐based cathodes.4FigureElectrochemical performance and stability of an anode‐supported single cell with columnar PCO cathodes. a) Cross‐sectional SEM image of a Ni‐YSZ anode‐supported single cell with a 2‐µm‐thick YSZ electrolyte layer and a 2‐µm‐thick columnar PCO20 cathode. b) Bright‐field TEM image of the contact between the PCO20 film and Pt CC. c) Typical I–V and I–P curves and d) impedance spectra of the single cell measured at T = 500–600 °C in wet H2 at the anode and air at the cathode. Note that I denotes the current density, V is the cell voltage, and P is the power density. e) Temperature dependence of ASR values corresponding to the total, polarization and ohmic resistance, respectively. f) Short‐term stability test of the single cell for 10 h measured at 500 °C under a constant current density of 0.3 A cm−2.When seeking to demonstrate robust high‐performance SOFCs at intermediate temperatures, several important cathode prerequisites need to be considered: 1) need for adequate levels of MIEC properties to ensure access to the entire oxide surface for the oxygen exchange process to proceed, 2) offer excellent catalytic activities towards ORR, and 3) exhibit high thermal/chemical stability. While the Co/Sr‐based perovskite oxides suffer from criteria 3, the fluorite PCO material suffers instead from low electronic conductivity (criteria 1) that limits current densities. While there have been many attempts over the years to address the chemical instability of Co/Sr‐based cathodes, for example, by modifying the composition and/or surface treatments,[17,23,60,61] the development and optimization of PCO as an electrode remain in their infancy. For the practical application of fuel cells with PCO, most importantly, the electrode has to be designed in such a way as to minimize the ohmic loss induced by the flow of low‐mobility electrons through the PCO layer. For this, first, it is essential to have sufficient Pr concentrations to support the higher level of electronic conductivity. Second, the electrode microstructure should allow electrons transferred from the current collector to be quickly conducted throughout the PCO surface as needed, for example, vertically‐ordered structures with short diffusion lengths as used in this work. Lastly, the micro‐contact between current collectors and PCO structures where ORR can occur on the entire PCO surface is required for the realization of large‐scale SOFCs for high power output. This suggests that there still remains scope to optimize the viable cell design technically for high‐performance and durable fuel cells.Discussion and ConclusionIn this work, we successfully demonstrate the feasibility of PCO as a robust high‐performance cathode for fuel cells by utilizing nanostructured thin‐film‐based PCO electrodes. Given the electrical conductivity of PCO as compared to LSC64 (Table S4, Supporting Information), further optimization, however, is still required to successfully integrate PCO into practical cells for widespread application. Hayd et al., for example, systemically investigated the influence of the microstructure of nano‐scaled LSC thin film cathodes on electrochemical performance by varying processing parameters such as processing temperature, heating rate, and annealing time.[62] As a result, a substantial increase in active surface area was achieved by control of grain size at the nanometric scale (≈17 nm) coupled with high porosity (45%), successfully leading to exceptionally low polarization resistance (0.023 Ω cm2 at 600 °C). Furthermore, the authors pointed out that the choice of the microstructure of the current collector could be optimized to minimize ohmic losses associated with low lateral electronic transport, an issue of concern in PCO. A popular alternative approach for optimizing current collection is the utilization of composite cathode structures. For example, in conventional SOFCs, (La, Sr)MnO3‐δ(LSM)‐YSZ composite structures are extensively used as cathodes given the high electronic and ionic conductivities of LSM and YSZ, respectively. By replacing YSZ, a pure ionic conductor, with MIEC PCO, it can be expected that, for example, a PCO/LSM composite would be much more active given that the ORR would not be limited to only the triple phase boundaries between LSM and YSZ but across the full PCO surface. With this in mind, further microstructural optimization of PCO can be expected in the future to achieve acceptable electrode performance in more conventional electrode configurations.From a scale‐up point of view, the PLD fabrication technique is not a representative method for scaling up this columnar electrode layer. However, columnar structures can be also fabricated via two representative methods for scaling up the electrode layers, which are magnetron sputtering[63] and chemical vapor deposition (CVD).[64] For example, Lee et al. successfully demonstrated an exceptional peak power density of 2.5 W cm−2 at 650 °C by employing columnar structures of the La0.6Sr0.4Co0.2Fe0.8O2.95‐YSZ cathode by a sputtering process.[63] In addition, Oh et al. also fabricated porous columnar thin films of undoped ceria with nano‐scaled grain size by metal‐organic CVD.[64] With these findings, suggested columnar structures, in themselves, should not restrict scale‐up, rather providing others with a guideline to further optimize fabrication processes.Furthermore, from the standpoint of cost and availability, cobalt (Co) and its strategic role presently being played in lithium battery technology today, driven by very rapid growth in electric vehicles, is likely to suffer rapid cost increases and supply disruptions in the not‐too‐distant future. In this regard, the development of Co‐free SOFC cathode material is strongly warranted for next‐generation fuel cells. By comparison, the price of the majority species in PCO, cerium oxide is approximately 25 times lower in cost than that of Co.[65] On the other hand, praseodymium oxide is presently comparable in cost to Co.[66] Based on our findings, while 40% Pr substitution for Ce gives the optimum performance, performance saturates above 10% substitution, offering substantial potential cost savings with little loss in performance.In summary, chemically stable and Co/Sr‐free fluorite‐based mixed conducting PCO was demonstrated, for the first time, to show exceptional intermediate temperature fuel cell performance, importantly coupled with much extended long‐term electrode durability. Electrical conductivity transient profiles confirmed that the kchem increases with Pr concentration, exhibiting a value of 3 × 10−5 cm s−1 for PCO20, comparable to the state‐of‐the‐art perovskite LSC64. The Pr ionization energy is proposed to serve as a key factor impacting the surface exchange kinetics, by impacting electron transfer from the PCO surface to the oxygen molecular adsorbates. Further, the vertically‐ordered, high surface area, columnar PCO structures, not only deliver outstanding electrode activity with an ASR lower than 0.1 Ω cm2 at 600 °C, but also a 15‐fold lower degradation rate of ≈0.2% h‐1 for 330 h compared to that of LSC64. Imposing peak power densities of 0.92 and 0.60 W cm−2 at 600 and 550 °C, respectively, were also achieved using a Ni‐YSZ anode‐supported single cell. Investigating and optimizing PCO as a SOFC cathode, therefore, serves as a promising vehicle for developing durable and high‐performance SOFCs that operate at intermediate temperatures.Experimental SectionFabrication of Dense and Vertically‐Ordered Columnar PCO Thin FilmsDense epitaxial PCO thin films with a thickness of ≈1 µm were grown on single‐crystal Al2O3 (0001) substrates (10 × 10 × 0.5 mm3, MTI Corporation) by PLD, operated with a KrF 248 nm excimer laser emitting at 248 nm (Coherent COMPex Pro 205) at an energy level of 300 mJ with a repetition rate of 10 Hz. PCO targets with different Pr concentrations were prepared by a combined EDTA‐citrate complexing method.[38] During the deposition step, the temperature and the working pressure were 650 °C and 10 mTorr O2, respectively. After the deposition process, the films were annealed at the same temperature under 1 Torr O2 for 20 min to ensure more complete oxidation of the films.Vertically‐ordered columnar PCO thin films were grown on both sides of single‐crystal YSZ (001) substrates (8 mol%, 10 × 10 × 0.5 mm3, MTI Corporation) by PLD (300 mJ, 10 Hz) with the same targets as indicated above. The deposition temperature was 600 °C and the working pressure was 100 and 250 mTorr O2. The film thickness varied from 0.5 to 2.5 µm. With a columnar LSC64 film as a reference for a comparison of the performance and stability, a dense Gd0.1Ce0.9O1.95 (GDC) buffer layer was initially deposited on both sides of the YSZ substrate by PLD at 700 °C under 10 mTorr O2 to prevent any unwanted reaction of YSZ with LSC64. The 1‐µm thick columnar LSC64 film was subsequently deposited on both sides of the GDC/YSZ substrate at 700 °C under 300 mTorr O2.Physical and Chemical CharacterizationThe microstructures of the deposited films were characterized using SEM (Hitachi S‐4800) and cross‐sectional scanning transmission electron microscopy (STEM, Titan cubed G2 60–300, FEI Co.). HR‐XRD (X'pert‐PRO MRD, PANalytical) measurements were taken for both in‐plane and out‐of‐plane reflections of the deposited PCO films using Kα (Cu) radiation (45 kV, 40 mA). The chemical composition of the PCO films was analyzed by an ICP‐MS and XRF. AFM with a Bruker (Innova) device in tapping mode and XPS (K‐alpha, Thermo VG Scientific) were used to investigate the surface roughness and the Si impurities of the films, respectively. The UV–vis spectrum was examined using a UV–vis spectrophotometer (Cary‐300, VARIAN) and UPS with a source energy of 21.21 eV was used to measure the energy difference between the conduction band and the Fermi level of the films.Electrical Conductivity Relaxation MeasurementsECR measurement was conducted by using two‐probe electrodes. With the aid of thin films used in this work, their geometric factor makes the lateral resistance through the films significantly large so that the contact resistance at electrodes should be negligible. This allows the measured resistance to be strongly dependent on the film, not at electrodes (e.g., contact resistance). Two Pt electrodes (200 nm thick and 1 mm width) as current collectors were prepared by DC magnetron sputtering (DC power of 100 W and Ar working pressure of 5 mTorr) with the aid of a shadow mask. The measurements were conducted in an alumina tube at temperatures of 600–650 °C with pO2 steps of 0.21 to 1 atm delivered using a four‐way valve via mass‐flow controllers (MFCs, Fujikin). In‐plane conductivity, which reflects the oxygen content in the PCO thin films, was monitored by measuring the voltage for every 0.5 s to acquire a reliable signal by means of chronopotentiometry (CP, VSP‐300, Biologic) until the sample adapted to a new equilibrium state. The normalized conductivity as a function of time was fitted using the first‐order Equation (1) to calculate the kchem.Electrochemical Measurements of PCO Symmetric CellsThe electrochemical performance of PCO symmetric cells was investigated by AC impedance spectroscopy (ACIS, VSP‐300, Biologic). Two types of current collectors were used: Sputtered metals (Pt and Au) and Ag paste. The cells were placed inside of a continuous‐flow alumina tube for the ACIS test under gas mixtures of O2 and Ar flowed through digital MFCs. Impedance spectra were obtained at temperatures ranging from 400 to 650 °C during the pO2 steps between 0.05 and 1 atm with AC perturbation of 20 mV in a frequency range of 2 MHz to 4 mHz. The performance stability of PCO and LSC64 symmetric cells was assessed for approximately 330 and 200 h, respectively, at 550 °C in dry and wet air atmospheres.Anode‐Supported Single Cell Fabrication and Electrochemical CharacterizationA vertically‐ordered columnar PCO20 cathode (2‐µm thick) was deposited by PLD on a commercial anode‐supported fuel cell with a thin YSZ (2‐µm thick) electrolyte (Elcogen Co.HC400B). The detailed deposition conditions for the PCO cathode are described above. A sputtered Pt current collector was then applied to the PCO cathode. A 2‐µm thick LSC64 cathode layer as a reference for comparison with a PCO‐based single cell was deposited on the GDC/YSZ substrate at 700 °C under 200 mTorr O2. The effective area of the cathode was 1 × 1 cm2. The cell testing configuration for the optimum current collection consisted of a metallic interconnect with a modified rib design and Ni‐foams and Au meshes.[67] Humidified hydrogen (3% H2O–97% H2) and air were applied as the fuel and oxidant to the anode and cathode, respectively. Impedance spectra were measured under open‐circuit voltage in a frequency range of 1 MHz to 0.1 Hz with an AC amplitude of 50 mV. The typical I–V and I–P curves of the single cells were recorded at operating temperatures varying from 600 to 500 °C at intervals of 50 °C using an Iviumstat electrochemical analyzer (Iviumstat, Ivium Technologies). The stability of the voltage was monitored under a constant current density of 0.3 A cm−2 at 500 °C for 10 h.AcknowledgementsThis work was supported by a National Research Foundation of Korea (NRF) grant funded by the Korean government (MSIT) (2021M3H4A1A01002695). J.‐W.S. and D.H.K. also appreciate financial support from Korea Institute of Energy Technology Evaluation and Planning (KETEP), the Ministry of Trade, Industry & Energy, Republic of Korea (No. 20213030030040). H.L.T. acknowledges support from U.S. Department of Energy (DOE), National Energy Technology Laboratory (NETL), Office of Fossil Energy under Award no. DE‐FE0031668.Conflict of InterestThe authors declare no conflict of interest.Author ContributionsH.G.S. and W.J. conceived the idea and designed the experimental protocol. H.G.S. performed the overall sample preparations, characterizations, electrical conductivity relaxation, and electrochemical measurements. D.H.K. and J.‐W.S. performed the fuel cell tests and helped to interpret the results. J.S. helped in collecting the conductivity relaxation data and assisted in the interpretation of the results. S.J.J. assisted with the collection of the thin film characterization data. J.K. assisted in the sample preparation. W.J. supervised the work and provided guidance throughout the project. H.L.T. assisted with the interpretation of the results. H.G.S., J.‐W.S., and W.J. wrote the manuscript with inputs from all co‐authors. All co‐authors contributed by discussing the results and all helped to revise the manuscript.Data Availability StatementResearch data are not shared.B. C. H. Steele, A. Heinzel, 2001, 414, 345.N. P. Brandon, S. Skinner, B. C. H. Steele, Annu. Rev. Mater. Res. 2003, 33, 183.Z. Shao, S. M. Halle, Nature 2004, 431, 170.H.‐I. Ji, J.‐H. Lee, J.‐W. Son, K. J. Yoon, S. Yang, B.‐K. Kim, J. Korean Ceram. Soc. 2020, 57, 480.S. Im, J.‐H. Lee, H.‐I. Ji, J. Korean Ceram. Soc. 2021, 58, 351.A. Evans, A. Bieberle‐Hütter, J. L. M. Rupp, L. J. Gauckler, J. Power Sources 2009, 194, 119.D. Beckel, A. Bieberle‐Hütter, A. Harvey, A. Infortuna, U. P. Muecke, M. Prestat, J. L. M. Rupp, L. J. Gauckler, J. Power Sources 2007, 173, 325.J. An, J. H. Shim, Y.‐B. Kim, J. S. Park, W. Lee, T. M. Gür, F. B. Prinz, MRS Bull. 2014, 39, 798.C.‐W. Kwon, J.‐W. Son, J.‐H. Lee, H.‐M. Kim, H.‐W. Lee, K.‐B. Kim, Adv. Funct. Mater. 2011, 21, 1154.M. Tsuchiya, B. K. Lai, S. Ramanathan, Nat. Nanotechnol. 2011, 6, 282.S. B. Adler, Chem. Rev. 2004, 104, 4791.A. Egger, E. Bucher, M. Yang, W. Sitte, Solid State Ionics 2012, 225, 55.Y. Chen, Y. Choi, S. Yoo, Y. Ding, R. Yan, K. Pei, C. Qu, L. Zhang, I. Chang, B. Zhao, Y. Zhang, H. Chen, Y. Chen, C. Yang, B. deGlee, R. Murphy, J. Liu, M. Liu, Joule 2018, 2, 938.E. Bucher, A. Egger, P. Ried, W. Sitte, P. Holtappels, Solid State Ionics 2008, 179, 1032.C. Xia, W. Rauch, F. Chen, M. Liu, Solid State Ionics 2002, 149, 11.S. W. Baek, J. H. Kim, J. Bae, Solid State Ionics 2008, 179, 1570.S.‐L. Zhang, H. Wang, M. Y. Lu, A.‐P. Zhang, L. V Mogni, Q. Liu, C.‐X. Li, C.‐J. Li, S. A. Barnett, Energy Environ. Sci. 2018, 11, 1870.Y. Zhang, B. Chen, D. Guan, M. Xu, R. Ran, M. Ni, W. Zhou, R. O'Hayre, Z. Shao, Nature 2021, 591, 246.S. Choi, S. Park, J. Shin, G. Kim, J. Mater. Chem. A 2015, 3, 6088.Y. Chen, S. Yoo, Y. Choi, J. H. Kim, Y. Ding, K. Pei, R. Murphy, Y. Zhang, B. Zhao, W. Zhang, H. Chen, Y. Chen, W. Yuan, C. Yang, M. Liu, Energy Environ. Sci. 2018, 11, 2458.Y. Chen, Y. Bu, Y. Zhang, R. Yan, D. Ding, B. Zhao, S. Yoo, D. Dang, R. Hu, C. Yang, M. Liu, Adv. Energy Mater. 2017, 7, 1601890.F. Prado, T. Armstrong, A. Caneiro, A. Manthiram, J. Electrochem. Soc. 2001, 148, J7.J. Wang, K. Y. Lam, M. Saccoccio, Y. Gao, D. Chen, F. Ciucci, J. Power Sources 2016, 324, 224.S. Hou, J. A. Alonso, J. B. Goodenough, J. Power Sources 2010, 195, 280.F. Wang, Q. Zhou, T. He, G. Li, H. Ding, J. Power Sources 2010, 195, 3772.H. Hayashi, M. Kanoh, C. J. Quan, H. Inaba, S. Wang, M. Dokiya, H. Tagawa, Solid State Ionics 2000, 132, 227.H. Hayashi, T. Saitou, N. Maruyama, H. Inaba, K. Kawamura, M. Mori, Solid State Ionics 2005, 176, 613.B. Koo, K. Kim, J. K. Kim, H. Kwon, J. W. Han, W. C. Jung, Joule 2018, 2, 1476.B. Koo, H. Kwon, Y. Kim, H. G. Seo, J. W. Han, W. Jung, Energy Environ. Sci. 2018, 11, 71.Z. Cai, M. Kubicek, J. Fleig, B. Yildiz, Chem. Mater. 2012, 24, 1116.W. Jung, H. L. Tuller, Energy Environ. Sci. 2012, 5, 5370.C. C. Wang, M. Gholizadeh, B. Hou, X. Fan, RSC Adv. 2021, 11, 7.E. Bucher, W. Sitte, Solid State Ionics 2011, 192, 480.A. F. Staerz, H. G. Seo, T. Defferriere, H. L. Tuller, J. Mater. Chem. A 2022, 10, 2618.C. Nicollet, C. Toparli, G. F. Harrington, T. Defferriere, B. Yildiz, H. L. Tuller, Nat. Catal. 2020, 3, 913.D. Chen, Z. Guan, D. Zhang, L. Trotochaud, E. Crumlin, S. Nemsak, H. Bluhm, H. L. Tuller, W. C. Chueh, Nat. Catal. 2020, 3, 116.D. Chen, S. R. Bishop, H. L. Tuller, J. Electroceram. 2012, 28, 62.H. G. Seo, Y. Choi, W. C. Jung, Adv. Energy Mater. 2018, 8, 1703647.M. Schaube, R. Merkle, J. Maier, J. Mater. Chem. A 2019, 7, 21854.H. Kim, H. G. Seo, Y. Choi, D.‐K. Lim, W. Jung, J. Mater. Chem. A 2020, 8, 14491.D. Chen, S. R. Bishop, H. L. Tuller, Adv. Funct. Mater. 2013, 23, 2168.S. R. Bishop, T. S. Stefanik, H. L. Tuller, Phys. Chem. Chem. Phys. 2011, 13, 10165.D. P. Fagg, V. V Kharton, A. Shaula, I. P. Marozau, J. R. Frade, Solid State Ionics 2005, 176, 1723.C. Lenser, F. Gunkel, Y. J. Sohn, N. H. Menzler, Solid State Ionics 2018, 314, 204.R. Chockalingam, A. K. Ganguli, S. Basu, J. Power Sources 2014, 250, 80.L. Zhao, N. H. Perry, T. Daio, K. Sasaki, S. R. Bishop, Chem. Mater. 2015, 27, 3065.C. B. Gopal, S. M. Haile, J. Mater. Chem. A 2014, 2, 2405.Y. Yin, S. Fu, S. Zhou, Y. Song, L. Li, M. Zhang, J. Wang, P. Mariyappan, S. M. Alshehri, T. Ahamad, Y. Yamauchi, Electron. Mater. Lett. 2020, 16, 224.J. J. Kim, S. R. Bishop, N. Thompson, Y. Kuru, H. L. Tuller, Solid State Ionics 2012, 225, 198.K. Schmale, M. Grünebaum, M. Janssen, S. Baumann, F. Schulze‐Küppers, H.‐D. Wiemhöfer, Phys. Status Solidi 2011, 248, 314.W. Jung, H. L. Tuller, Adv. Energy Mater. 2011, 1, 1184.T. S. Stefanik, Electrical Properties and Defect Structure of Praseodymium‐Cerium Oxide Solid Solutions, Ph. D., Massachusetts Institute of Technology, Cambridge, MA 2004.A. Infortuna, A. S. Harvey, L. J. Gauckler, Adv. Funct. Mater. 2008, 18, 127.J.‐H. Park, W.‐S. Hong, K. J. Yoon, J.‐H. Lee, H.‐W. Lee, J.‐W. Son, J. Electrochem. Soc. 2014, 161, F16.D. Beckel, U. P. Muecke, T. Gyger, G. Florey, A. Infortuna, L. J. Gauckler, Solid State Ionics 2007, 178, 407.J. Yoon, R. Araujo, N. Grunbaum, L. Baqué, A. Serquis, A. Caneiro, X. Zhang, H. Wang, Appl. Surf. Sci. 2007, 254, 266.M. Shang, J. Tong, R. O'Hayre, RSC Adv. 2013, 3, 15769.C. Duan, D. Hook, Y. Chen, J. Tong, R. O'Hayre, Energy Environ. Sci. 2017, 10, 176.M. Li, M. Zhao, F. Li, W. Zhou, V. K. Peterson, X. Xu, Z. Shao, I. Gentle, Z. Zhu, Nat. Commun. 2017, 8, 1.B. Koo, J. Seo, J. K. Kim, W. Jung, J. Mater. Chem. A 2020, 8, 13763.N. Tsvetkov, Q. Lu, L. Sun, E. J. Crumlin, B. Yildiz, Nat. Mater. 2016, 15, 1010.J. Hayd, L. Dieterle, U. Guntow, D. Gerthsen, E. Ivers‐Tiffée, J. Power Sources 2011, 196, 7263.Y. H. Lee, H. Ren, E. A. Wu, E. E. Fullerton, Y. S. Meng, N. Q. Minh, Nano Lett. 2020, 20, 2943.T.‐S. Oh, D. A. Boyd, D. G. Goodwin, S. M. Haile, Phys. Chem. Chem. Phys. 2013, 15, 2466.U.S. Geological Survey, Mineral Commodity Summaries 2022, U.S. Geological Survey, Reston, VA, 2022.Statista, Praseodymium Oxide Price Worldwide from 2009 to 2020 with a Forecast from 2021 to 2030 (in U.S. Dollars per Metric Ton), Stormcrow, 2021.H. S. Noh, J. Hwang, K. Yoon, B. K. Kim, H. W. Lee, J. H. Lee, J. W. Son, J. Power Sources 2013, 230, 109.

Journal

Advanced Energy MaterialsWiley

Published: Nov 1, 2022

Keywords: (Pr,Ce)O 2; Co/Sr‐free mixed conducting oxides; nanocolumnar structures; oxygen reduction reaction; solid oxide fuel cells

References