Ordering and Phase Control in Epitaxial Double-Perovskite Catalysts

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Ordering and phase control in epitaxial doubleperovskite catalysts for oxygen evolution reaction Felix Gunkel, Lei Jin, David N. Mueller, Clemens Hausner, Daniel Bick, ChunLin Jia, Theodor Schneller, Ilia Valov, Rainer Waser, and Regina Dittmann ACS Catal., Just Accepted Manuscript • DOI: 10.1021/acscatal.7b02036 • Publication Date (Web): 07 Sep 2017 Downloaded from http://pubs.acs.org on September 7, 2017

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Ordering and phase control in epitaxial double-perovskite catalysts for oxygen evolution reaction Felix Gunkel,∗,† Lei Jin,‡,§ David N. Mueller,¶ Clemens Hausner,† Daniel Bick,† Chun-Lin Jia,‡,§ Theodor Schneller,† Ilia Valov,¶ Rainer Waser,¶ and Regina Dittmann¶ †Institute of Electronic Materials (IWE2), RWTH Aachen University, 52074 Aachen, Germany ‡Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, Forschungszentrum Juelich GmbH, 52425 Juelich, Germany ¶Peter Gruenberg Institute and JARA-FIT, Forschungszentrum Juelich GmbH, 52425 Juelich, Germany §Peter Gruenberg Institute and JARA-FIT, Forschungszentrum Juelich GmbH, 52425 Juelich, Germany E-mail: [email protected]

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Abstract The complex oxide compound praseodymium barium cobalt oxide (PBCO) is an efficient catalyst for the oxygen evolution reaction (OER) during electrochemical water splitting, with an activity that is mainly ascribed to PBCO’s inherent atomic structure and band alignment. Here, we report on epitaxial PBCO thin films showing electrocatalytic properties, with current densities of up to 10 mA/cm2 at 1.8V vs. RHE. Dense PBCO thin films are synthesized in a disordered perovskite phase as well as in an coherently oxygen vacancy-ordered (double)perovskite phase, in which oxygen vacancies are incorporated in every second CoO2−δ atomic plane along the out-of-plane direction. The transition from disordered to ordered growth occurs temperature-controlled during the growth process and can be directly monitored in-situ by means of reflection high-energy electron diffraction. The epitaxial fabrication process allows the control of structure and phase of the oxide catalysts, providing model systems for exploring structure-property relations and atomistic processes of catalysis during OER. For all structural compositions, we demonstrate remarkably similar catalytic properties, indicating a negligible effect of the structural bulk phase on OER catalysis. Rational design routes for perovskites catalysts derived merely from bulk properties should therefore be met with suspicion.


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1 Introduction Chemical storage of energy is a key challenge for evolving from fossil large facility energy production to local energy production by means of renewable energy sources. Hydrogen is considered one of the most important media for such energy storage. 1–6 Sustainable and efficient production of hydrogen is essential in this process. One particular focus of research in this field is electrolysis of water, in which hydrogen is produced directly by splitting water molecules into its components oxygen and hydrogen under applied electrochemical stimulus. The efficiency of electrolysis of water is thus an essential challenge for realizing sustainable energy storage in renewable energy concepts. While the basic principle of electrolysis and catalysis has been known for decades, 7–10 the atomistic processes at active oxide surfaces catalyzing the water-splitting reaction are still highly debated. 11–17 One of the greatest challenges is the identification and determination of the decisive physical descriptors that promote electrocatalytic performance. In particular, in the search for new materials catalyzing the oxygen evolution reaction (OER), reflecting the anodic half-reaction of the water splitting process, different bulk descriptors of electronic (band structure, orbital hybridization) and structural (crystal structure, surface morphology, active sites, etc.) 11–13,15,18–22 nature have been proposed. The separation of these descriptors is often difficult, as typically more than one of these parameters is being changed among different samples at the same time and in an entangled manner. Complementing the search for descriptors, chemical stability of catalysts upon electrochemical treatments is another central aspect in realizing effective and sustainable applications. 23–25 Besides noble metals and binary metal oxides such as IrO2 or RuO2 , 14,26–29 (double)- perovskites are being studied as active materials for OER 5,11,24,30–34 as well as for fuel cell applications 35 in recent years. This is 1) because of the wide variety of available compositions and stoichiometry in this class of materials and 2) because of their specific surface structure which provides potentially active sites for adsorption of reactants 18,20,36 and promotes oxygen exchange between active electrode and electrolyte. 37 Stoichiometry variations 38–40 , tailored defect structures, 41,42 and space charge formation 43–46 yield additional degrees of freedom to control electric 3 ACS Paragon Plus Environment

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Figure 1: (a) Schematic of OER taking at the active oxide surface. (b)-(e) Possible structural appearance of PBCO (double-)perovskites. (b) Fully disordered (Pr0.5 Ba0.5 ) CoO3−δ ; (c) A-site disordered (Pr0.5 Ba0.5 )2 Co2 O5+δ with oxygen vacancy within every other nominal CoO2 plane. (d) A-site ordered PrBaCo2 O5+δ with oxygen vacancy in nominal PrO atomic plane; (e) A-site ordered PrBaCo2 O5.5+δ with oxygen vacancy in every other nominal CoO2 atomic plane.

properties, band alignments and catalytic activity in complex oxides on the nanoscale, but at same time, render complex oxide catalysis an elaborate scientific problem, particularly in ternary or even quaternary systems. The concentration, distribution, and function of defects — most importantly oxygen vacancies — is a particular complexity in the field. Complex oxide catalysts are synthesized by chemical routes resulting in powders or polycrystalline (and even porous) thin films 24,30,31 , maximizing the active surface area utilized in applications. While being feasible for technological applications, such samples sometimes make it difficult to disclose the decisive material parameters responsible for catalytic activity in a systematic manner. Here, we report on electrochemically active, epitaxial praseodymium barium cobalt oxide (PBCO) thin films, which show significant current densities upon electrochemical water splitting, and at the same time dense and defined structural properties (Fig. 1a). The thin films are obtained with phase control of the ordered and disordered bulk phase and nearly atomic precision, allowing for a direct connection of the catalysts’ structural properties on the atomic level and the resulting electrochemical activity probed in-operation. As we demonstrate, the disordered phase, in which oxygen vacancies are expected to be statistically distributed, shows similar catalytic activity as the ordered PBCO phase in which oxygen vacancies are incorporated preferentially in every second CoO2−δ atomic plane along the out-of-plane direction and are coherently ordered. 4 ACS Paragon Plus Environment

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Atomic ordering of A-site cations such as discussed in the literature 30,47–49 does not occur in either phase, revealing new insights into the atomic structure of PBCO as ordered double-perovskite compounds. The absence of a clear (bulk) structure-property relation in OER catalysis implies a generally skeptical perspective on bulk arguments often used to derive rational design rules for perovskite oxide catalysts.

Atomic structure of PBCO PBCO may crystallize in the simplest case in a perovskite lattice, with Pr3+ and Ba2+ sharing the A-site in the lattice 50 and mixed-valence Co2+/3+/4+ cations on the B-site, surrounded by O2− octahedra (referred to as (Pr0.5 ,Ba0.5 )CoO3−δ , see Fig. 1b). Oxides generally contain oxygen vacancies, indicated by δ , whereas concentration and distribution of these defects can be controlled by thermodynamic treatments and/or chemical doping. Due to the low oxidation state of Co cations, cobaltites contain a particularly large concentration of oxygen vacancies, potentially yielding oxygen non-stoichiometry of several tens of percent (i.e. δ may take values up to 0.8, while maintaining cubic symmetry of the lattice 51 ). As a result, oxygen vacancies can show coherent ordering rather than statistical distribution within the lattice. 18,32,52 In this case, oxygen vacancies are in various compounds found to be located within every other CoO2−δ atomic plane 53–55 (in the PBCO case referred to as (Pr0.5 Ba0.5 )2 Co2 O5+δ , Fig. 1c), which in the extreme case may lead to the formation of a Brownmillerite structure such as reported for SrCoO2.5 . 56–58 ) Mixed A-site cobaltites, such as PBCO, may however exhibit even more complex atomic configurations, this is an ordered double-perovskite structure, in which the A-site cations order coherently, too (see e.g. Figs. 1d,e), mainly owing to the different ionic radii of the cation species. 47–49 In this case, structural oxygen vacancy formation has been observed to occur in the PrO-atomic planes (Figs. 1d, referred to as PrBaCo2 O5+δ ). Finally, a combination of A-site ordering and oxygen vacancy formation in every other CoO2−δ atomic plane (Fig. 1e, referred to as PrBaCo2 O5.5+δ ) may be possible. Hence, PBCO may occur in various structure and atomic configurations. 5 ACS Paragon Plus Environment

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Associated with its specific atomic structure and stoichiometry, PBCO supposedly provides a nearly ideal eg band filling as well as energetic position of the O2p-band center for catalyzing OER. 12,30 Accordingly, powder samples and polycrystalline thin films have proven sufficiently low over-potential as well as good chemical stability and long-term performance of PBCO during OER. 24,30,31 At elevated temperatures, also fast oxygen exchange kinetics have been reported for PBCO 36,59 leading to considerable efforts of using PBCO in solid oxide fuel cells 60 and oxygen sensors. 32 Given the complex atomic structure of mixed A-site cobaltites, however, the validation of structure-property relations frequently being used in rational design rules for catalysts requires careful processing and a detailed analysis. Judgment of electrocatalytic performance should be accompanied by careful determination of the atomic structure of the catalyst, in order to derive a relation between material chemistry and electrocatalytic performance.

2 Results PBCO thin films with thicknesses of 100 − 120 nm were grown by RHEED-controlled pulsed laser deposition (PLD) on (100) SrTiO3 (STO) single crystal substrates at an oxygen pressure of 0.1 torr, a laser fluence of 3.5 Jcm−2 , and a repetition rate of 5 Hz using a ceramic PBCO target. The growth temperature was varied between 650 ◦ C and 950 ◦ C (heater temperature). All thin films initially grow in a defined layer-by-layer growth mode as revealed by RHEED intensity oscillation observed during the PLD growth (Fig. 2a). After a few monolayers, the RHEED intensity gradually drops until the end of the growth process. Figs. 2b-d illustrate the evolution of the RHEED intensity during the initial growth phase for various growth temperatures. At 650 ◦ C, we observe a weak single-frequency oscillation, corresponding to layer-by-layer growth of single perovskite unit cells. At increased growth temperature (Fig. 2c), much clearer RHEED intensity oscillations evolve, whereas every second oscillation damps out over time (indicated by arrows pointing upwards), eventually resulting in a doubling of the oscillation period (arrows pointing downwards). This behavior becomes even more pronounced at a growth temperature of 950 ◦ C,


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for which only weak intermediate oscillations are observed, while the entire RHEED characteristic is governed by the doubled growth period. This is an unusual behavior, given that the growth rate in PLD is determined by constant parameters such as laser fluence, background pressure and frequency 38,61 , as far as volatility of the species is negligible. 62 Therefore, the unusual change in the growth period hints towards a transition of the growth entities, i.e., the doubling of the material required for the formation of a single monolayer of PBCO. As we will elaborate in more detail below, this transition in the growth entity reflects a controlled transition from PBCO-nucleation in a common perovskite lattice (short period, Fig. 1b) to a nucleation in an oxygen-vacancy ordered (quasi-)double-perovskite structure (doubled period, Fig. 1c). As shown in Fig. 2e for 100 nm thick layers, all films show a defined and dense surface morphology. The mean roughness of these films is typically of the order of RMS < 2 nm and constant among the samples. The flat surface of the epitaxial PLD-thin films allows a precise estimation of the real active surface area in electrochemical model experiments, which differs only of order 0.1% from the geometric surface of the sample (assuming negligible effects of active-area blocking due to gas bubble formation formation reported and characterized in the literature for catalyst 7 ACS Paragon Plus Environment

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layers with differing surface roughness micrometer-scale 63 ). The intended cation stoichiometry of the PBCO thin films was confirmed by inductively coupled plasma mass spectrometry (ICP-MS) revealing Pr/Ba/Co-ratio of 1/0.99/2.08. Crystal structure and phase of the thin films were analyzed by X-ray diffraction (XRD) displayed in Fig. 3. As evident for all growth temperatures, clear thin film peaks appear at slightly larger angles as compared to the STO substrate (Figs. 3a and b), indicating the formation of c˚ Furthermore, all oriented perovskite-type PBCO thin films with a lattice spacing about 3.87 A. PBCO layers show a 4-fold in-plane symmetry as well as rocking curves close to single crystal substrate quality (FWHM values of the rocking curves are below 0.045◦ ; representative ϕ -scans


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and rocking curves are shown in the supplementary information SI1). Interestingly, at increased growth temperature of 900◦ C and higher, additional peaks appear in the XRD scans. The most pronounced peak is observed at about 11.5◦ corresponding to an out-of-plane lattice spacing of ˚ (Figs. 3c), i.e., a coherent doubling of the perovskite unit cell, reflected by appearance 7.75 A of half-integer-indexed peaks ((00 12 ), (00 32 ), and (00 52 )). (Note, that the occurrence of half-order peaks in XRD implies atomic distortions of the unit cell, e.g. atomic displacements or shifts, as will be discussed below.) The thin films grow in a strained manner, as indicated by the (013)reciprocal space maps (RSM) delivering STO substrate peak and PBCO thin film peak at the same in-plane lattice constant (Figs. 3d). A similar result is obtained for low-temperature (left) and high-temperature (right) grown samples. However, only for samples grown at high temperature, we also observe finite intensity for the (01 52 ) peak corroborating the appearance of coherent order and a doubled unit cell (Fig. 3e). In contrast, low-temperature-grown samples do not show any intensity in the same region, confirming the absence of coherent order in these thin films. PBCO thin films can thus be stabilized both in a disordered phase (low growth temperatures, cf. Fig. 1b) and in an ordered phase (elevated growth temperature). The atomic nature of this ordering process (cf. Figs. 1c-e), however, cannot be resolved based on RHEED and XRD and requires atomic scale analysis. Fig. 4a shows the microstructure of a PBCO thin film grown at 650◦ C (here only 20 nm thick) obtained by high-resolution scanning transmission electron microscopy (HRSTEM), revealing a few unit cell thick coherent layer formed at the interface to the STO substrate, consistent with the observation of RHEED-intensity oscillations. At larger thickness, a disordered grain-like microstructure is revealed, similar to the one reported in the literature 64,65 , while a general c-axis orientation is maintained. At increased growth temperature of 850 ◦ C (Fig. 4b), a more defined microstructure is obtained throughout the entire film thickness. Strikingly, and corroborating the foregoing structural analysis, various regions show a fringe contrast (red arrows). As evident from Fourier transformed image (FT, taken from the area marked by the green box), the this area shows tetragonal periodicity as expected for a strained perovskite lattice, but with additional doubling of the out-of-plane lattice parameter, as indicated by extra intensity observed at half


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Figure 4: HRSTEM analysis for samples grown at 650◦ C (a), 850◦ C (b) and 950◦ C (c). The enlarged image in (a) illustrates the disordered perovskite phase obtained at low growth temperature. (b) In some regions, fringe contrast is observed at 850◦ C, indicating double-spaced atomic ordering, confirmed by the appearance of half-order peaks after Fourier transformation (top right). (c) Films grown at at 950◦ C show coherent ordering within the entire thin films evident in the overview image as well as in the magnified inset (red arrows indicate ordered low-intensity CoO2−δ planes). White arrows indicate Ruddlesden-Popper-like stacking faults. The intensity line scan displayed in the bottom panel reveals decreased intensity in every second B-site plane, and enlarged interatomic distances between every second A-site plane (line scan taken between the black lines indicated in the top panel).

distance (green arrows). At further increased growth temperature (950◦ C, Fig. 4c), wide regions of this ordered perovskite structure appear, in line with the increasing intensity and appearance of superlattice peaks in the corresponding XRD patterns. Meanwhile, planar defects can be observed, as indicated by the white open arrows. Although these defects are still present even at the highest growth temperature, their density is found to be much lower (Fig. 4c and inset) than at intermediate growth temperature. A closer investigation of these planar defect structures reveals a local rock salt structure with multiple A-site atomic planes associated with Ruddlesden-Popper-like local phases. The fringe contrast observed for high-temperature-grown samples is characterized by a


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lowered relative intensity in every second CoO2−δ atomic plane, indicating that oxygen vacancies are formed mainly in the CoO2−δ atomic planes (Figs. 1c,e), rather than in AO planes (Fig. 1d). The intensity profile (bottom panel of Fig. 4c) taken perpendicular to the ordering direction (indicated by the black arrow) reveals not only the lowered intensity of every other B-site atomic plane (red arrows), but also an enlarged interplanar spacing between the A-site cations in the vicinity of the oxygen deficient CoO2−δ planes (green bars), yielding the lowered symmetry of the unit cell responsible for the superlattice peaks observed in XRD. The finding of oxygen vacancies being located in the B-site atomic planes of PBCO contradicts some literature reports on PBCO powders (typically treated at temperatures above 1000 ◦ C), 30,49 but is consistent with reports on SrCoO3−δ thin films 56 , (La,Ba)CoO3−δ 32 , (Y,Sr)CoO3−δ 53 , and (La,Sr)MnO3−δ 54,55 . Both, limited kinetics during thin film growth as well as epitaxial strain (which is absent in powders) may explain these differences. In order to investigate A-site cation ordering in the thin films, chemically sensitive analysis was carried out by atomic-scale electron dispersive x-ray spectroscopy (EDXS) for all samples. Fig. 5 displays typical results obtained for the sample grown at 950◦ C. While the atomic positions of Pr and Ba A-site sublattice (Figs. 5b,c) and Co B-site sublattice (Fig. 5c) show spacially complementary intensity, no obvious contrast is observed among the two A-site cations, which show intensity contrast in similar atomic columns. Hence, there is no evidence for A-site cation ordering (along the viewing direction) from the EDXS experiment (Note, that EDXS revealed a similar absence of A-site ordering for all samples investigated here).


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Based on structural analysis, we obtained a systematic transition from fully disordered to an oxygen-vacancy ordered crystal phase solely controlled by the growth temperature of the thin films. At low temperature, no ordering is observed. At intermediate growth temperature, a mixed phase involving ordered and disordered phases is established. At the highest growth temperature, a fully oxygen vacancy-ordered (double-)persovkite PBCO thin films is obtained, in which oxygen vacancies are formed within the CoO2−δ plane, while A-site cation ordering is not evident. As shown in Fig. 6, despite of the clear differences in the structural phase, the electrical conductivity of PBCO does not scale with growth temperature apart from a general scatter. The measurements instead suggest that the electronic properties of PBCO are independent from the actual structural phase and atomic structure of PBCO. (Note that due to the highly insulating STO substrates the electrical conductivity of PBCO is measured in absence of any chemical admixtures 30,33 or metallic current drains 24,31 , often used to increase the overall conductivity in powder catalysts.) Other vacancy-ordered perovskites such as SrCoO3−δ 56 show a pronounced metal-toinsulator transition upon phase transformation. However, in the case of PBCO, the conductivity seems to be solely dictated by the mixed valence of A-site cations, yielding a highly doped and well-conducting material in both ordered and disordered phase. Additionally, while ordering of


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oxygen vacancies occurs only at elevated growth temperatures, a significant amount of oxygen vacancies seems to be present also in the disordered phase, yielding comparably high electrical conductivity values. The in-plane conductivity of several hundred S/cm up to 2000 S/cm is remarkably high for perovskites and among the largest values reported for such compounds. 6 Despite of the extremely high conductivity values, however, a 100 nm thick PBCO film possesses a total sheet resistance in the range of hundred to several hundred Ohms (right axis in Fig. 6). For electrochemical characterization, 100 nm thick Pt contact pads are sputtered on the edges of the 10 × 10 mm2 epitaxial PBCO/STO samples, providing Ohmic contact to the active PBCO layer. The Pt pads are then connected to the potentiostat leads using copper tape. In the center


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of the sample, an o-ring of diameter 5 mm seals the active PBCO/electrolyte (1M KOH) interface from the electrical leads and prevents any contact of the sputtered Pt contact pads to the electrolyte (inset of Fig. 7a). Within the electrochemical cell we employ a Hg/HgO reference electrode (Radiometer analytical) and a large-area Pt counter electrode, such as frequently used in the characterization of OER catalysts. 14,24,30,31,33,37 The use of platinum enables a fast cathodic reaction on the counter electrode side (hydrogen evolution reaction, HER) of the cell. In this way, it is ensured that the total water splitting reaction is limited by the anodic oxygen evolution reaction taking place at the working electrode under test, i.e. the epitaxial PBCO thin film electrodes. Metal dissolution from the Pt counter electrodes (observed in reverse polarization, applied to test HER or oxygen reduction catalysts) is not expected under the applied bias during OER. 5,66–68 Figure 7a shows the raw results of cyclic voltammetry (CV) experiments obtained from the epitaxial PBCO thin film samples. Here, we plot the total current density as a function of the applied potential, E vs. RHE, referenced to reversible hydrogen electrode potential. The catalytic PBCO thin films on STO generate pronounced non-zero current density reaching several milliamps per square-centimeter at potentials above about 1.7 V, clearly demonstrating oxygen evolution reaction taking place at the epitaxial PBCO surfaces. For all samples, we performed multiple CV-sweeps (up to 100 cycles). Apart from a slight drift, the samples show stable CV-characteristics. Above the OER-threshold, the current density increases with different slopes, corresponding to resistance values of a few hundreds of Ohms, i.e., a similar order of magnitude as the sheet resistance values found in van-der-Pauw geometry. Hence, at (high) potentials above 1.8 Vvs.RHE, the thin films’ resistance seems to reflect the major limit for the current density observed during OER. In fact, due to the sample geometry, the PBCO thin films act like a bottleneck for electron transport in the circuit, as the entire charge has to drain through the PBCO thin film (see inset of Fig. 7a). Therefore, the different slopes obtained for differently grown samples indicate different total resistance of the PBCO thin films (which scatters among the samples), rather than inherent differences in catalytic performance. (Note however that hydrophilicity and surface blocking by gas bubble formation on the nanoscale may differ from sample to sample, too.) Electrochemical impedance spectroscopy


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(EIS) reveals consistent total Ohmic serial resistances, RΩ , in the circuit. EIS measurements were carried out for each sample at frequencies between 1 Hz and 100 kHz. RΩ was determined from the high-frequency intersect with the real axis in the Nyquist-plot (cf. Supplementary Information SI2 for further details). Typically, RΩ is attributed to the electrolyte resistance, in particular when using powder catalysts and/or current collecting admixtures. For the samples investigated in this study, however, RΩ is a complex quantity with contributions of electrolyte resistance, but also (and dominantly) of the PBCO thin film. While a direct comparison of RΩ obtained from EIS and RS obtained from van-der-Pauw is difficult due to the much more complex geometry applied in the EIS experiment, the similarity of their values may imply that the current density at large potentials is mainly limited by the drain of electrons being transferred into the PBCO layer upon OER and cutting off the mass turnover, rather than by the rate of the actual surface reaction (nevertheless, the surface reaction leads to dominant contributions in the overpotential, as will be discussed below). After IR-correction, the CV-characteristics reflect the actual catalytic properties of the individual surfaces of the PBCO thin films, yielding current density of up to 10 mA/cm2 at 1.8 V vs RHE. As shown in Fig. 7b, the CV-characteristics of the different samples now almost overlap indicating that the inherent catalytic activity of low temperature and high temperature grown samples is rather similar. In fact, despite of the enormous differences in microstructure and crystal phase, all epitaxial (100) PBCO thin films show very comparable CV-behavior, suggesting that the crystal phase and in particular order and disorder in these films has only limited effect on the catalytic performance. Instead, the results emphasize the importance of the thin films surface and the catalyst/electrolyte interface, respectively. The structural ordering of PBCO on (100) STO is epitaxially stabilized in parallel to the thin film/substrate interface and thus in parallel to the sample’s surface. Therefore, the surface structure exposed to the electrolyte during OER — corresponding to the projection along the (100) direction of the two phases (i.e., top view of Figs. 1b,c) — is similar in both cases. As discussed in the literature, the surface structure of the thin film may furthermore differ from the bulk phase, due to solid/liquid interaction, local undercoordination and leaching effects. 14,24,30


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While the details of this interface process remains to be elucidated, the similar behavior of ordered and disordered phase revealed in this study demonstrates that catalytic activity of these materials is not solely determined electronically, e.g. by the bulk band structure or Co-O-hybridization (which is expected to clearly differ for the two phases in response to the modified coordination of Co ions within the lattice). This fact readily reflects the complexity of OER catalysis and complicates the search for suitable catalysts descriptors. Mere bulk properties of the PBCO thin films cannot be used to derive rational design rules for catalyst fabrication.

3 Conclusion In conclusion, we have demonstrated the phase control in epitaxial PBCO thin films, synthesized in a disordered and an ordered (quasi) double-perovskite crystal structure. In the ordered phase, ordering arises from the formation of oxygen vacancies within every second cobalt oxide atomic plane, in absence of A-site cation ordering, reflecting a new insight into the atomic structure of PBCO as an oxygen-vacancy ordered (but A-site cation disordered) perovskite. In fact, in this atomic structure it may even be questionable if can be called a double-perovskite at all, as this generally implies non-degeneracy of A-sites within the lattice. The nucleation of the ordered phase is determined by the growth temperature, thus controlled by thermodynamics and growth kinetics. The epitaxial PBCO thin films are highly conducting with electrical conductivities in the 1000 S/cm-range and show good catalytic activity for OER in alkaline media [...]. As demonstrated, the electrochemical activity is independent of the structural and atomistic phase of the PBCO thin films, illustrating the importance of the catalyst’s surface chemistry, while generally questioning rational design rules for OER perovskite catalysts based on bulk arguments. Such as demonstrated here, epitaxial complex oxide thin films can combine atomic phase control and catalytic activity close to large-surface-area samples, thus providing an ideal model to study the details of the oxide-electrolyte interface during OER in the future.


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