Insights into the Surface Reactivity of Cermet and Perovskite

Subscriber access provided by UNIVERSITY OF CONNECTICUT

Article

Insights into the surface reactivity of cermet and perovskite electrodes in oxidizing, reducing and humid environments Fotios Paloukis, Kaliopi Papazisi, Thierry Dintzer, Vasiliki Papaefthimiou, Viktoriia A. Saveleva, Stella P. Balomenou, Dimitrios Tsiplakides, Fabrice Bournel, Jean-Jacques Gallet, and Spyridon Zafeiratos ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b05721 • Publication Date (Web): 06 Jul 2017 Downloaded from http://pubs.acs.org on July 7, 2017

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a free service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are accessible to all readers and citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

ACS Applied Materials & Interfaces is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Applied Materials & Interfaces

Insights into the surface reactivity of cermet and perovskite electrodes in oxidizing, reducing and humid environments Fotios Paloukisa, Kalliopi M. Papazisib, Thierry Dintzera, Vasiliki Papaefthimioua, Viktoriia A. Savelevaa, Stella P. Balomenoub, Dimitrios Tsiplakidesb,c, Fabrice Bourneld,e, Jean Jacques Galletd,e and Spyridon Zafeiratosa,* a

Institut de Chimie et Procédés pour l'Energie, l'Environnement et la Santé, UMR 7515 du

CNRS-UdS 25 Rue Becquerel, 67087 Strasbourg, France b

Chemical Process and Energy Resources Institute /CERTH, 6th km Charilaou-Thermi Rd.,

57001 Thessaloniki, Greece c

Department of Chemistry, Aristotle University of Thessaloniki, 54124 Thessaloniki, Greece

d

Laboratoire de Chimie Physique-Matière et Rayonnement, Sorbonne Universités – UPMC Univ.

Paris 06 – CNRS, 4 place Jussieu, 75005 Paris, France e

Synchrotron-Soleil, L'orme des Merisiers, Saint Aubin – BP48 91192 Gif-sur-Yvette Cedex,

France KEYWORDS: Electrolysis, solid oxide electrochemical cells, NAP-XPS, perovskites, cermets.

ACS Paragon Plus Environment

1


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 2 of 43

ABSTRACT: Understanding the surface chemistry of electrode materials under gas environments is important in order to control their performance during electrochemical and catalytic applications. This work compares the surface reactivity of Ni/YSZ and La0.75Sr0.25Cr0.9Fe0.1O3 which are commonly used types of electrodes in solid oxide electrochemical devices. In situ synchrotron-based near ambient pressure photoemission and absorption spectroscopy experiments, assisted by theoretical spectra simulations and combined with microscopy and electrochemical measurements, are used to monitor the effect of the gas atmosphere on the chemical state, the morphology and the electrical conductivity of the electrodes. It is shown that the surface of both electrode types readjusts fast to the reactive gas atmosphere and their surface composition is notably modified. In the case of Ni/YSZ this is followed by evident changes in the oxidation state of nickel, while for La0.75Sr0.25Cr0.9Fe0.1O3 a fine adjustment of Cr valence and strong Sr segregation is observed. An important difference between the two electrodes is their capacity to maintain adsorbed hydroxyl groups on their surface, which is expected to be critical for the electrocatalytic properties of the materials. The insights gained from the surface analysis may serve as a paradigm for understanding the effect of the gas environment on the electrochemical performance and the electrical conductivity of the electrodes.

INTRODUCTION Solid oxide electrochemical cells (SOC) are high temperature electrochemical devices that can convert fuels to electricity (fuel-cell mode) or alternatively to store electricity as chemical fuels (electrolysis mode). Currently the most commonly used fuel electrode material in SOC devices is Ni/Yttria Stabilized Zirconia cermet (Ni/YSZ hereafter), mainly due to its low cost and good mechanical strength.1 Nevertheless, even though Ni/YSZ-based SOCs have demonstrated


2

Page 3 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


excellent initial performance, there are various reports that stress some disadvantages such as low tolerance to gas impurities (mainly sulphur) and poor redox stability, causing nickel agglomeration and cell deactivation after prolonged operation.1–3 In response to these failures, novel materials are continuously being processed as candidates for SOC fuel electrodes. A possible alternative for Ni/YSZ cermets are electrodes made of perovskite-based materials. Several studies have manifested improved performance of perovskite fuel electrodes employed in fuel-cells4,5 and electrolyzers.6–8 The effect of the gas atmosphere on the electrode stability is particularly important in applications where the electrodes should operate both under oxidizing and reducing atmospheres, as for example in symmetrical SOCs,9,9 or when fuel flexibility is required. In addition, redox treatments of cermet SOC electrodes are often related to the reduction of the interface between the electrode, the electrolyte and the gas (triple phase boundary, TPB), the formation of cracks and in general, to the performance deterioration.11–13 In a porous gas electrode, the electrochemical reaction takes place most preferably at the TPB where the gaseous fuel phase, the ion conducting phase (electrolyte), and the electron conducting phase, are in contact, while in the case of mixed conducting materials the reaction zone is largely expanded and electrochemical reactions take place at the surface; therefore, optimizing the TPB content and understanding the nature of the surface play a prominent role and provide excellent opportunities for SOCs performance improvement.14–16 In many cases, changes under gas treatments, especially in the microstructure, are irreversible and therefore it is possible to detect them by post-treatment analysis of the electrodes.13 However, several recent studies have shown that the electrode/electrolyte and electrode/gas interfaces are dynamic and respond to changes of the gas phase,17,18 temperature19,20 and the applied potential.21 Examples include the polarization-driven


3


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 4 of 43

exsolution of B-site transition-metal cations from the perovskite lattice22,23 and the formation of an optimum surface composition of Ni-gadolinium-doped ceria (Ni-GDC) cermet electrodes for maximum cell currents.24 Therefore, modifications related to the surface oxidation state and interdiffusion within the outer surface layers are difficult to observe in the post-treatment analysis and require in situ electrode characterization with methods of high surface sensitivity. Despite the fact that it was developed only very recently, Near Ambient Pressure X-ray Photoelectron Spectroscopy (NAP-XPS) has been proven as one of the primary techniques to study electrochemical devices under working conditions.15,25 The surface modification of perovskite-type materials under reactive gas environments has been extensively studied by NAPXPS,26–32 while studies of Ni-based cermet electrodes are relatively scarce.16,17,24,33 Most of the high-temperature studies involve model SOC electrodes developed in order to conform with the requirements of the method and facilitate the electrochemical operation.34 Although these papers gave valuable insights, the comparison of the experimental findings obtained is difficult due to differences in the experimental conditions (e.g. temperature, pressure, X-ray radiation, etc.) and the materials preparation.35 In this article, we compare the surface chemistry of the two most commonly used types of SOC electrodes (cermet and perovskite) at intermediate temperature and at gas-phase environments relevant for both SOC modes, fuel cell and electrolysis cell. Ni/YSZ was chosen because it is the state-of-the-art electrode in most practical SOC applications, while doped perovskites of the LSCM type (i.e. La-Sr-Cr-Metal) exhibit inherent redox stability and are potential candidates for use in symmetrical and reversible SOCs.36 Given the sensitivity of the surface state and composition to the gas environment, the study was focused on in situ NAPXPS, NEXAFS and electrochemical characterization, combined with the morphological and


4

Page 5 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


structural characteristics measured on post-treated samples. Special care has been taken to ensure that the treatment conditions are kept identical by performing parallel experiments of the two electrodes. Several similarities, but also significant differences on the reactivity of the two electrode materials relative to gases are found, which can help to elucidate the impact of the surface on the performance of these electrodes and hopefully to develop more robust electrodes for SOC applications. EXPERIMENTAL SECTION Sample preparation The two electrodes were fabricated using a typical preparation procedure for solid oxide button cells.37 In particular, commercial NiO-YSZ cermet (Fuel Cell Materials, 66% by weight NiO, 34% by weight (Y2O3)0.08(ZrO2)0.92) and laboratory synthesized5 La0.75Sr0.25Cr0.9Fe0.1O3 perovskite (LSCrF hereafter) were formed as inks (solids content 70 %wt.) and applied by means of screen printing (STV, mesh 40 µm) on an Yttria-Stabilized-Zirconia (YSZ) pellet (Kerafol, 8YSZ) with a thickness of 150 µm. Between YSZ electrolyte and LSCrF electrode a CGO (Gd0.20Ce0.80O1.95, Fuel Cell Materials) interlayer was applied in order to prevent forming the non-conducting La2Zr2O738 after prolonged exposure in high operating temperatures. The mean Ni/YSZ and LSCrF electrodes’ thickness was about 40 µm as estimated by cross-sectional SEM images. Before being introduced in the NAP-XPS chamber, the electrodes were thermally treated in air at 1350°C for 3 h in the case of NiO-YSZ and at 1200°C for 3 h for the LSCrF electrode. Synchrotron-based experiments Near Ambient-Pressure X-ray photoemission and absorption spectroscopies (NAP-XPS and NEXAFS) were carried out at the NAP-XPS end station of the University Pierre et Marie Curie set on the TEMPO beamline of the SOLEIL synchrotron radiation facility in France. The setup


5


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 6 of 43

was equipped with a SPECS Phoibos 150-NAP electron analyzer and the incident photon beam has a diameter of 100 µm. The sample was heated using a button heater and the temperature was measured with a K-type thermocouple. The gas flows into the analysis chamber were regulated by leak valves and monitored by a quadrupole mass spectrometer. For the measurements, the Ni/YSZ and LSCrF samples were mounted together on a ceramic heater and fixed in position with two stainless steel clamps (supporting information, fig. S1). The temperature was monitored by a thermocouple attached on top of the samples. NAP-XPS spectra were collected using selected photon energies, so that the obtained photoelectrons have the same kinetic energy (190 eV) and thus identical sample information depth (about 2 nm). In case of the Sr 3d and O 1s photoelectron peaks, spectra were recorded with an additional higher incident photon energy, which gave photoelectrons of 590 eV and an estimated information depth of about 5 nm. The presented NAP-XPS spectra were acquired using a 20 eV pass energy. The binding energy scale is referenced to the Fermi edge of the grounded sample which was measured subsequently to the core levels, using the same photon energy. In the absence of states at the Fermi edge (e.g. oxidized Ni/YSZ), its position was estimated indirectly from the valence band features. For more details on valence band region spectra please refer to supporting information S2. The peak areas were estimated after subtraction of a Shirley background and the quantitative calculations were performed based on the peak areas of Ni 2p, Zr 3d, Y 3d (for Ni/YSZ) and La 3d, Sr 3d, Fe 2p and Cr 2p, (for LSCrF) photoelectron peaks, taking into account the photo-ionization crosssections dependence of the atomic subshells.39 It is worth mentioning that the absolute values of the elemental atomic concentrations obtained by this approach might not be very accurate, however it is safe to compare relative values between different measurements. For the Sr 3d peak fitting two set of Sr 3d doublets were used with fixed width, peak area ratio and spin orbit


6

Page 7 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


splitting. In case of the O1s peaks a mixed Gaussian/ Lorentzian (70/30%) line shape was used with the peak position and width constrained at different ranges for each component. Finally, the Ni 2p peak in H2O was deconvoluted using reference Ni 2p peak shapes recorded in fully reduced and oxidized nickel samples. The samples were measured at total pressure of 3.5 (±0.5) mbar under 5 different atmospheres applied in the consecutive order: O2, H2, H2/H2O (1:1), H2/H2O (1:3), H2O, O2 and ultra-high vacuum conditions (UHV, ca. 5x10-8 mbar). The sequence of the experiments is shown in detail in supporting information, fig. S3. The measurements were performed as follows: first, the sample was annealed at 500 °C (± 5 °C) in the indicated atmosphere for 30 min and subsequently the temperature was lowered at 300 °C and NAP-XPS and NEXAFS measurements were recorded keeping the same gas atmosphere. This procedure was considered necessary in order to avoid overheating of the analyzer nozzle which was in close proximity to the sample. The carbon contaminates on the surface were monitored by recording the C 1s spectra under all applied conditions. After the initial oxygen treatment, adventitious carbon is removed from the surface and the C 1s signal becomes negligible. However, the C 1s peak increases again in subsequent H2 ambient up to 6 %, while in the following H2O/H2 mixture carbon is removed. Please note that the behavior is identical for both electrodes which indicates that the origin of the carbon signal is probably due to carbon residuals in the chamber, or in the gas lines that are deposited on the electrodes, rather than bulk diluted carbon that segregates on the surface of the electrodes. Photon beam damage effects were tested by shifting the analysis spot to a new sample position and the lack of differences in the spectra manifests the stability of the samples under the photon beam in the time frame of the experiment.


7


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 8 of 43

The NEXAFS spectra were recorded in the Auger Electron Yield (AEY-NEXAFS) mode, enhanced by additional electrons created by ionization of the gas phase above the sample. The estimated information depth of NEXAFS measurements is about 5 nm.40 The experimental Cr Ledge spectra were simulated using the so-called charge-transfer multiplet (CTM) approach.41,42 The calculations have been carried out using the CTM4XAS523 program43 and literature values were used as guides for the Slater integrals (Fdd, Fpd, Gpd), the spin-orbit splitting parameter (SO), the difference between the core hole potential Upd and the 3d-3d repulsion energy Udd, as well as for the hopping parameters (eg and t2g).44 The octahedral symmetry is chosen for the calculations of both Cr3+ and Cr4+.45 In particular for the simulation of octahedrally coordinated Cr3+ the parameters used were : (Fdd, Fpd, Gpd)=0.9, SO=1.04, Upd-Udd =1, eg =1.35 and t2g=0.65. In case of octahedrally coordinated Cr4+ the parameters used were: (Fdd, Fpd, Gpd)=0.9, SO=0.95, Upd-Udd =3, eg =1 and t2g=0.5. Simulations of the NAP-XPS spectra were performed by the Simulation of Electron Spectra for Surface Analysis (SESSA) software, Version 2.0 (National Institute of Standards and Technology, Gaithersburg, MD). Homogeneous distribution, layer-by-layer and spherical particles on substrate morphologies were modeled. The surface composition of the LSCrF electrode was calculated based on the employed models. Morphology, structure and electrochemical characterization The surface morphology was inspected by scanning electron microscopy (SEM) using a Zeiss GeminiSEM 500 microscope. Energy-dispersive X-ray spectroscopy (EDXS) mapping was combined with SEM images to resolve the different elements on the surface of the electrodes. An electron beam energy of 12 and 6 keV was used for LSCrF and Ni/YSZ, respectively. The


8

Page 9 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


powder X-ray diffraction (XRD) experiments were carried out with a Bruker D8 diffractometer. The XRD pattern was recorded in the 10–80 (2θ) range with a scan step of 0.02. Electrodes’ ohmic resistance (Rohm) was calculated by AC-impedance measurements at 500 0C under different gas atmospheres. The stabilization time before the measurements was 30 min. A three-electrode cell configuration was used, where Ni/YSZ or LSCrF was the working electrode in each case, La1−xSrxMnO3 (LSM) perovskite was applied as counter electrode and platinum paste as reference electrode. Counter and reference electrodes were co-located on the same side of the YSZ substrate, while the counter LSM electrode was symmetrical and located opposite to the working electrode, and at a distance of 3 mm from the Pt reference electrode. The AC impedance spectra were recorded with a potentiostat/galvanostat (µAutolab-type III) at open circuit potential (OCP). The frequency range was from 100 kHz to 10 mHz with a signal amplitude of 10 mV. The ohmic resistance component (Rohm) was obtained from the intercept of the high frequency part of the impedance spectrum with the real axis, while the polarization resistance of the electrodes can be determined by subtracting the previous ohmic contribution to the total resistance obtained from the intercept with the real axis at low frequencies. The contribution of the other resistances (fine mesh grids and lead wires) is very small since Pt is a high conductive material.46 The three-electrode configuration allowed the measurement of the impedance at the working electrode interface without the influence of the counter electrode.47 The experiments were carried out in a variable-pressure 0.6 L reactor attached to a laboratorybased ultrahigh-vacuum (UHV) setup.48 The gas inlet and gas detection of this system were analogous to those in the synchrotron based setup. RESULTS AND DISCUSSION The surface oxidation state in O2, H2, H2O and H2/H2O mixtures.


9


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 10 of 43

The valence state at the surface of Ni/YSZ and LSCrF samples in equilibrium with various gases is discussed first. Figures 1a and b show the Zr 3d and Y 3d NAP-XPS spectra of Ni/YSZ electrode in O2. The Zr 3d5/2 and Y 3d5/2 peaks at 182.3 and 157.1 eV respectively, are in good agreement with previous reports of zirconium and yttrium atoms in yttria-stabilized zirconia.49,50 The binding energy (BE) and the shape of these peaks remain identical in H2 and H2O atmospheres (not shown) indicating a stable valence state of YSZ under the conditions employed. On the contrary, the oxidation state of nickel is modified by the gas atmosphere as shown by the Ni 2p peaks presented in figure 1c. In particular, in O2 the main Ni 2p3/2 peak at 855.5 eV is accompanied by an intense satellite around 862 eV as expected for oxidized Ni2+.24,51 When the gas atmosphere is switched to H2 the Ni 2p3/2 peak shifts to 852.7 eV, while the intensity of the accompanying satellite peak decreases, indicating the reduction of nickel oxide to metallic nickel (Ni0).24,51 The Ni 2p peak is not modified in H2/H2O mixtures (for brevity only 1:1 H2/H2O ratio is shown). However, when H2 gas is no longer supplied and the Ni/YSZ electrode is exposed and measured under pure H2O atmosphere, nickel is almost completely oxidized back to Ni2+ with a minor contribution of metallic Ni0 (ca. 7%), as shown by the deconvoluted Ni 2p spectrum in figure 1c. Since the binding energies of oxide (NiO) and hydroxide (Ni(OH)2 or NiOOH) nickel species are very close52 one cannot safely assign the Ni 2p peak to one of the above species. However, the O 1s spectra show only a minor contribution of hydroxyl species (see figure 7 below), indicating that in H2O atmosphere nickel is mainly oxidized to nickel oxide (e.g. NiO) and not to hydroxide. This means that most probably the mechanism of nickel oxidation in presence of steam involves fast decomposition of adsorbed H2O on nickel to form NiO and not Ni(OH)2.


10

Page 11 of 43

a) Zr 3d

2+

c) Ni 2p

Ni

0

Ni

H2O

Intensity / a.u.

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


O2 188

186

184

182

180

H2O/H2

b) Y 3d

H2

O2 O2 162

160

158

156

Binding Energy / eV

154

880

870

860

850

Binding Energy / eV

Figure 1 . (a) Zr 3d (hν=375 eV), (b) Y 3d (hν=350 eV) and (c) Ni 2p (hν=1050 eV) NAP-XPS spectra of the Ni/YSZ electrode recorded at 300 °C in 3.5 mbar of the indicated gas atmosphere. Prior to spectra acquisition the electrode was annealed at 500 °C for 30 min in the indicated gas atmosphere. The NAP-XPS spectra of LSCrF electrode under identical conditions are presented in figure 2. Similar to Ni/YSZ electrode, the La 3d, Fe 2p, Cr 2p peaks have the same shape and binding energy in all studied environments and therefore for brevity only spectra recorded in O2 are shown, while spectra in reducing condition can be found in supporting information S4. The main La 3d5/2 peak (fig. 2a) at 834.6 eV is accompanied by a strong shake up satellite shifted by about 3.9 eV.53 The intensity and energy separation of the satellite relative to the main peak are known to be highly sensitive to the ligand atoms around La.53 The spectrum (peak BE and satellite peak


11


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 12 of 43

separation) in figure 2a is not affected by the gas treatments or the excitation photon energy and is in good agreement with previous reports for La3+ ions in the perovskite structure.54 The Fe 2p peak shown in fig. 2b has low signal to noise ratio due to the relatively low amount of iron doping in the perovskite. Although this complicates its analysis, the characteristic BE at 710.8 eV suggests that Fe is most probably in the Fe3+ valence.55 The Cr 2p3/2 peak (fig. 2c) appears at 576.2 eV and designates trivalent (Cr3+) cations.56 The Auger electron yield NEXAFS spectra of the Fe and Cr L3,2-edges are more informative to the iron and chromium oxidation state, due to an improved signal to noise ratio as compared to NAP-XPS. In addition, NEXAFS spectra are sensitive to the coordination geometry of the cations, therefore they can reveal changes on the perovskite surface structure. The Fe L3,2-edge spectrum shown in figure 2d consists of two components at about 710 and 722 eV, corresponding to the Fe L3 and L2 edges respectively, due to the Fe 2p spin orbit splitting.57 The shape of the Fe L3,2-edge remains unaffected under all examined atmospheres (for brevity H2O/H2 mixtures are not included in the graph) and is alike with previously reported spectra of octahedral Fe3+ in α-Fe2O3.55,57 The NEXAFS findings not only support the NAP-XPS results, but also suggest that the Fe cations (B-site) are located at the body center of the cubic perovskite structure (ABO3) and maintain this position independently of the gas atmosphere.


12

Page 13 of 43

d) Fe L3,2-edge

O2

864 858 852 846 840 834 828

Intensity / a.u.

a) La 3d

O2 H2O

b) Fe 2p3/2

H2 705

Intensity / a.u.

710

O2

715

e) Exp. Cr L3-edge

725

720

725

Photon Energy / eV

720

715

710

c) Cr 2p

f) Theor. Cr L3-edge

705

Intensity / a.u.

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


3+

b

Cr + 4+ 0.3Cr

O2

c

3+

H2O

a

Cr

H2

O2

4+

Cr

Diff. O2-H2 595

590

585

580

575

Binding Energy / eV

570

576

580

584

576

580

584

Photon Energy / eV

Figure 2. (a) La 3d (hν=1030 eV), (b) Fe 2p (hν=900 eV), (c) Cr 2p (hν=770 eV) NAP-XPS spectra and (d) Fe L3,2-edge, (e) Cr L3-edge AEY-NEXAFS spectra of the LSCrF electrode. The lower curve is obtained after subtraction of the edge in H2 from the one in O2 after proper


13


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 14 of 43

normalization. All spectra were recorded at 300 °C in 3.5 mbar of the indicated gas atmosphere. Prior to spectra acquisition the electrode was annealed at 500 °C for 30 min in the indicated gas atmosphere. (f) Theoretically simulated Cr L3 edge spectra for octahedral (Oh) coordinated Cr3+ (top) and Cr4+ cations. The linear combination of Cr3++ 0.3Cr4+ spectra is presented at the top. The Cr L3-edge spectrum in Figure 2e (for the complete spectrum with both L3 and L2 edges please see supporting information S5) is dominated by a complex structure at about 578 eV. The line shape and the energy position of the Cr L3-edge resembles previously reported spectra of Cr2O3,58,59 which further confirms the deriving conclusions from NAP-XPS for the presence of Cr3+ species. However, in contrast with the Fe spectra which were unaffected by the gas atmosphere, the presented Cr L3 edges reveal some slight modifications in the features marked as a and c, which are clearly visible in the difference spectra obtained after proper normalization of the original spectra (bottom curve, fig. 2e). The peak shape of the difference curve is similar to previous NEXAFS spectra of Cr4+ ions,58 while it is clearly different for spectra reported for Cr6+ compounds.59 In order to interpret the modifications in the Cr NEXAFS spectra at different environments we simulated the Cr L3-edge of octahedrally coordinated Cr3+ and Cr4+ ions using charge transfer multiple calculations (CTM). As shown in figure 2f, a good accordance with the experimental spectra is given for Cr3+ octahedrally coordinated with a crystal field value of 10Dq=1.2 eV, and a charge transfer energy value ∆=7. In addition, the difference between the experimental curves in H2O and O2 ambient (bottom of fig. 2e) can be simulated by Cr4+ octahedrally coordinated with a crystal field value of 10Dq = 1.6 eV, and a charge transfer energy value ∆=1. As described previously,60 the 10Dq value is sensitive to the distance between Cr and O ions in the octahedral crystal, while ∆ is a measure of their interaction strength. In particular, the higher the


14

Page 15 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


10Dq value the lower the distance between Cr and O ions, while stronger Cr-O bonds will decrease the ∆ value. Of course, one should note that the provided 10Dq and ∆ parameters should be used only in comparative basis since they are determined by comparison between simulated and experimentally derived spectra and not by ab initio calculations. A linear combination of the theoretical curves of octahedral Cr3+ and Cr4+ cations in a (Cr3++ 0.3Cr4+) composition (top of fig. 2f) can describe the increase of a and c spectral features observed at the experimental results in O2. Overall, analysis of Cr L3-edge reveals that there is a significantly higher amount of Cr4+ cations in LSCrF structure in O2 atmosphere, as compared to the reducing environments. The presence of Cr4+ cations is not unexpected and is directly related to Sr-doping in the perovskite structure. In the ideal LSCrF perovskite structure (ABO3), Sr and La atoms occupy the hole which is created by eight Cr and Fe octahedra. During doping the trivalent La3+ is substituted by the divalent Sr2+ at the perovskite A-site. The introduced effective negative charge is compensated either by an increase in the valence of the B-site Cr3+ cations (electronic compensation) and/or by the formation of oxygen vacancies (ionic compensation), to maintain the electroneutrality.61,62 NEXAFS analysis confirms the presence of Cr4+ cations in LSCrF and suggests that their decrease in reducing atmospheres should be related to the modification of the Sr cations. Indeed, the NAP-XPS Sr 3d spectrum changes significantly with the gas atmosphere. Figure 3 shows the Sr 3d spectra recorded under O2, H2 and H2O atmospheres (for brevity, spectra of mixed H2/H2O are omitted). The presented Sr 3d spectra were obtained using two incident photon energies (325 and 725 eV) in order to distinguish depth-resolved spectral features.63 In O2 (figure 3a) two Sr 3d doublets at 131.6 and 133.0 eV can be identified in the spectrum. When the


15


gas is switched to H2 (figure 3b) the peak at 131.6 eV decreases considerably and the spectrum is dominated by the peak at 133.00 eV. In the mixed H2/H2O atmosphere the peak at 131.6 eV disappears (not shown) and remains absent even in H2O (figure 3c). b) Sr 3d in H2

a) Sr 3d in O2

surface

Intensity / a.u.

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 16 of 43

c) Sr 3d in H2O

sub-surface

hv=725 eV

hv=325 eV

138

136

134

132

Binding Energy / eV

130

138

136

134

132

130

Binding Energy / eV

138

136

134

132

130

Binding Energy / eV

Figure 3. Sr 3d NAP-XPS spectra recorded on LSCrF electrode using 725 eV (top) and 325 eV (bottom) incident photon energies in 3.5 mbar of (a) O2, (b) H2 and (c) H2O atmospheres. All spectra were recorded at 300 °C after the electrode was annealed at 500 °C for 30 min in the indicated gas atmosphere. The NAP-XPS measurements with the higher analysis depth (hv=725 eV) show an enhanced contribution of the component at 131.6 eV as compared to the lower depth (hv=325 eV). This is a clear indication that the Sr species at 131.6 eV are located in deeper layers (sub-surface) than those at 133 eV (surface or Srδ+). Formation of two Sr 3d components with distinct depth distribution has been reported in several previous XPS studies of perovskite electrodes.27,28,30,64 The low BE component (131.6 eV) is typically attributed to Sr species from the perovskite


16

Page 17 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


lattice, while the one at higher BE (133 eV) to surface Sr species,27,28,30,64 in accordance with the presented results. Here it is shown that in the presence of H2 the sub-surface component gradually disappears from the perovskite surface (the maximum analysis depth at the employed conditions is estimated at about 4.7 nm) and remains absent even in 100 % H2O atmosphere. The disappearance of the bulk Sr 3d component implies that all Sr cations (in the analysis depth of NAP-XPS) abandon their lattice position to segregate on the surface. It is interesting that when the sample is re-exposed to O2 at the end of the described treatment cycles, the Sr 3d spectrum returns to its initial shape (data not shown here). This indicates that the observed modifications are reversible and rapid in response to the gas phase changes. Surface composition in response to the gas atmosphere. NAP-XPS can be used to estimate the near-surface atomic composition of the electrodes (ca. 2-3 nm) and more importantly, to quantify in situ modifications induced by the different gas atmospheres. Figure 4a shows the evolution of the % atomic ratio of Ni at different gas environments as determined from the photoelectron peak areas. The reduction of nickel oxide due to exposure in H2 (shown in figure 1c) is followed by a decrease in the surface nickel concentration. This trend continues in H2/H2O mixtures, while it is reversible when the gas phase is switched back to O2 at the final oxidation step. The changes in the nickel atomic concentration suggest that the surface is rearranged and in particular, that in H2 and H2O atmospheres the concentration of nickel diminishes as compared to O2. This can be explained by diffusion of nickel into the YSZ particles and/or shrinkage of the nickel particles volume induced by nickel oxide reduction. Similar surface rearrangement has been reported for Ni/GDC electrodes.17,24 There, the electrodes were exposed to the same gases, but at one order of magnitude lower pressure. The consistency in the two cases despite the significant pressure difference, shows that


17


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 18 of 43

surface rearrangement is a general characteristic of nickel-based cermet electrodes and not an effect of the low gas pressure.24 It should be stressed that although the final O2 treatment indicates the reversibility of the surface composition in response to the gas atmosphere, figure 4a clearly shows that the calculated compositions in O2 atmosphere (initial and final) are significantly different. What we may observe here is the evolution of the Ni oxidation from the first to the final exposure in O2. One should keep in mind that in a standard Ni/YSZ SOC electrode, the first treatment step is always the reduction under H2 atmosphere, while in the reported experiments exposure in O2 has been the first step, followed by exposure in H2 (reduction). Our results point out that the surface composition under operation might be affected by the pre-treatment routine. Further experiments testing different pre-treatment conditions (gas type and temperature) as well as repeated redox cycles need to be performed, in order to provide insight into the effect of treatment procedures on the reversibility of the surface composition. The influence of the gas atmosphere is also clearly evident in the elemental composition of the LSCrF electrode shown in Figure 4b. In O2 atmosphere the surface of the LSCrF is dominated by La, while in H2 and H2O atmospheres La is gradually replaced by Sr. The concentrations of Cr and Fe are less affected by the gas atmosphere. Upon re-exposure to O2 the composition returns back to the initial one, revealing that surface reconstruction is fast and fully reversible. The rearrangement of the four perovskite cations in different atmospheres is better described by their relative ratio presented in supporting information, fig. S6, where it is shown that the ratio between Cr, Fe and La cations is stable and remains unaffected by the gas atmosphere. On the contrary, in hydrogen and humid environments Sr surface concentration gradually increases up to 25 times, while it rapidly returns to its initial value upon re-oxidation.


18

Page 19 of 43

100

a) Ni/YSZ NAP-XPS

% Ni % YSZ

% Atomic ratio

80

EDXS Ni

60 40 YSZ

20 0

O2 100 80

% Atomic ratio

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


H2

H2/H2O

H2O

O2(final) % La % Sr % Cr % Fe

b) LSCrF NAP-XPS

EDXS

60 Cr La

40 20

Sr Fe

0

O2

H2

H2/H2O

H2O

O2(final)

Figure 4. Bar chart showing the atomic concentration at the outer 2 nm of the surface as calculated in the basis of NAP-XPS peak areas normalized by the theoretical sensitivity factors (a) Ni/YSZ and (b) LSCrF electrodes. The atomic ratios in the bulk as estimated by the EDXS analysis are shown as dotted lines. Morphology and structure of reduced and oxidized electrodes. SEM combined with EDXS mapping was employed to investigate the morphology of the electrodes after exposure in H2 and the final O2/UHV treatments. Please note that the


19


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 20 of 43

measurements were performed ex-situ on sample quenched after H2 or O2 and UHV treatments and stored in air for about four days prior to the measurements. Although air exposure is expected to form an oxide layer on the surface, this will not induce significant morphology changes, since the surface mobility is limited at room temperature. Figures 5a and b show topview SEM micrographs of Ni/YSZ taken after reduction in H2. The electrode structure consists of bright nanoporous nickel particles mixed with dark-colored, smooth YSZ grains, as indicated by the EDXS mapping (figure 5c). Formation of nanopores in nickel particles upon reduction of Ni/YSZ cermets has been observed by environmental TEM and has been attributed to the volume contraction when the initially present NiO is reduced to Ni.65,66 As nickel oxide grows on the surface after the reoxidation and UHV annealing treatments (figure 5d, e), the nanopores into Ni particles close while their texture appears rougher. Furthermore there are no well-defined boundaries between the Ni and YSZ grains, which are evidently more overlapping as compared to the reduced sample. The SEM images are consistent with the NAP-XPS results of figure 4a which showed decrease of the nickel atomic ratio in reducing environments and increase upon reoxidation (final oxidation).


20

Page 21 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 5. a, b) Top view SEM micrographs of the Ni/YSZ electrode after annealing for 30 min at 500 °C in 3.5 mbar H2 c) SEM image with superimposed EDXS map (Zr: green; Ni: orange). e), d) Top view SEM micrographs of the Ni/YSZ electrode after consecutive treatments (of about 24 hours) performed during the NAP-XPS experiments which terminated with a cycle of O2 and UHV annealing. In the case of the LSCrF electrode, the grain particles after reduction appear relatively homogeneous without significant contrast between the particles in the image (figure 6a, b). In addition, the LSCrF particles have large, flat terraces and smooth edges. After the O2/UHV treatment (figure 6c) there is no considerable modification of the macroscopic surface morphology. The porous size remains, while the particles still exhibit clear facets. However, in the higher magnification image (figure 6e), it is evident that the surface of LSCrF particles is decorated by agglomerates which vary in size and shape, from ca. 20 nm round, to 50 -100 nm random-shaped particles.


21


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 22 of 43

Although it would be tempting to associate the particles observed in the SEM images with the strong surface enrichment of Srδ+ revealed by NAP-XPS this would be an oversimplification, while as discussed below the situation is more complex. The EDXS analysis is not conclusive over the composition of the small agglomerates over the large particles, since it shows a homogeneous EDXS distribution of all elements (figure 6d). The difficulty to distinguish differences in the surface composition on the LSMF particles by EDXS is justified by their relatively small thickness (estimated at about 50 nm from the SEM images) and the high information depth of EDXS analysis (up to 2000 nm). The X-ray diffraction patterns of Ni/YSZ and LSCrF electrodes after reducing and oxidative treatments are shown in supporting information, fig. S7. As expected, the cubic YSZ phase was always present in the Ni/YSZ electrode, while cubic Ni and monoclinic NiO phases were dominant in the reduced and oxidized samples, respectively. Accordingly, not only the surface, but also the bulk volume of the nickel particles, reacts under the employed low-pressure conditions. In the case of the LSCrF electrode, the XRD patterns correspond to a single phase with rhombohedral perovskite structure. The patterns of reduced and oxidized LSCrF electrodes appear identical, which indicates that there is not any prominent phase transition in the LSCrF structure and the modifications shown in fig. 4b are restricted to the surface. Previous studies indicated rhombohedral to orthorhombic phase transformation of the perovskite structure in La0.75Sr0.25Cr1-xMnxO3 electrodes upon hydrogen annealing at temperatures higher than 1000 °C.67 This supports the similarity of the XRD results of reduced and oxidized electrodes in our case, since the temperature was too low to affect the LSCrF structure.


22

Page 23 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 6. a, b) Top view SEM micrographs of the LSCrF electrode after annealing for 30 min at 500 °C in 3.5 mbar H2 c, e) Top view SEM micrographs of the LSCrF electrode after consecutive treatments (of about 24 hours) performed during the NAP-XPS experiments which terminated with a cycle of O2 and UHV annealing. d) SEM image with superimposed EDXS map (Sr: purple; La: green, Cr: peach Fe: orange). Adsorbed oxygen species. The O 1s peaks of the Ni/YSZ electrode exposed in various gas atmospheres are shown in figure 7a. The spectra are composed by a primary feature due to oxygen species on Ni/YSZ and a separate peak shifted at higher BEs due to the contribution of the gas atmosphere above the sample.68 In case of Ni/YSZ, two O 1s components at 530.5 and 531.8 eV are used to fit the overall peak. The main component around 530.5 eV has been assigned to a convolution of lattice


23


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 24 of 43

oxygen species from both NiO and YSZ.69 The peak at 531.8 eV has been attributed both to hydroxide species and defective sites within the oxide crystal.70 The intensity ratio between the two oxygen components is shown at the right part of figure 7a. It is evident that in the presence of humidity the contribution of the peak at 531.8 eV is increasing as compared to O2 and H2 atmospheres. The O 1s spectra of LSCrF electrode presented in figure 7b are similar to those previously recorded for LSF perovskite electrodes22 and are not affected by the type of the gas atmosphere. Peak deconvolution shows that in all cases the O 1s peak can be fitted by two O 1s components. The dominant O 1s component at 530.7 eV has been related to the surface oxygen layer of the perovskite lattice29 and oxygen species around the vacant oxygen sites (oxygen vacancies),71 but it is also in the energy range for lanthanum oxides53 and Sr2Co2O5 films.28 Given the relatively small deviation in the reported binding energies it is not possible to distinguish, on the basis of the O 1s peak, whether this component is due to La and Sr surface oxides or to oxygen from the perovskite lattice (including defect sites). However, the relatively high width of the peak at 530.7 eV (FWHM=1.79 eV) suggests that several oxygen species with similar BEs contribute to this O 1s component. The low-energy peak at 529.4 eV (FWHM=1.35 eV) is typically assigned to oxygen from the subsurface (bulk) perovskite lattice.28,29,71 It is interesting that the O 1s spectra remain apparently the same in oxidative, reducing and humid atmospheres (both H2/H2O and H2O) even if there is a significant modification between Sr and La surface concentration, as shown in figure 4b. This suggests that the oxide lattice structure, both surface and subsurface, is not significantly perturbated by the Sr and La cation interdiffusion induced by the gas environment.


24

Page 25 of 43

a) Ni/YSZ

b) LSCrF

530.5 eV 531.8 eV

530.7 eV 529.4 eV

H2O

H2O/H2

Intensity / a.u.

H2O/H2

2

H

2

H

O

O -g as

-g as

H2O

H2

O

2

-g as

-g as

H2

2

O

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


O2

O2

540 538 536 534 532 530 528

Binding Energy / eV

0,1

0,2

0,3

O1s(531.8)/O1s(530.5)

540 538 536 534 532 530 528

Binding Energy / eV

0,1

0,2

0,3

O1s(529.4)/O1s(530.7)

Figure 7. O 1s (hv=725 eV) NAP-XPS spectra recorded at 300 °C in 3.5 mbar of the indicated gas atmosphere for (a) Ni/YSZ and (b) LSCrF electrodes. Prior to spectra acquisition the electrode was annealed at 500 °C for 30 min in the indicated gas atmosphere. The graph at the right part shows the evolution of the O 1s peak components intensity ratio at each gas atmosphere. The oxygen gas-phase peak is split into two separate components due to the paramagnetic nature of the O2 molecule. Depth resolved O 1s spectra recorded in H2O are shown in the supporting information fig. S8. Although the direct comparison of the shown depth profiles with thin film data may not be straightforwardly possible due to the rough surface morphology of our electrodes, valuable insights about the nature of the oxygen species can be given by this analysis. In case of Ni/YSZ


25


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 26 of 43

the component at 531.8 eV is preferentially located over the oxygen species appearing at 530.5 eV. Accordingly, the component at 530.5 eV can be assigned to oxide lattice species (the peak resolution does not allow to distinguish if it originates from the YSZ or NiO lattice) and the peak at 531.8 eV to adsorbed oxygen species. Depth information, in combination with its enhancement with humidity (as shown in Figure 7a) and the BE position, provide a sound evidence that the peak at 531.8 eV is to a large extent due to adsorbed hydroxyl groups formed over the Ni/YSZ surface. It is interesting that the amount of the hydroxyl groups does not seem to be influenced from the steam partial pressure (almost the same amount in 50, 70 and 100 % steam), which indicates that the surface is saturated by hydroxyl-groups under the employed temperature and pressure conditions. In the case of LSCrF electrode the depth-dependent measurements (figure S8) show that the relative intensity of the two oxygen species is independent of the analysis depth. This allows us to conclude that both oxygen species are part of the perovskite lattice and/or the segregated oxide layers over it. Therefore, there is no evidence of adsorbed oxygen species on the LSCrF electrode even in H2O atmosphere (see figure 7b). This is well-correlated with the absence of significant O 1s signal between 531 and 533 eV which on previous NAP-XPS studies was attributed to adsorbed OH species.21,27 The difficulty of LSCrF electrodes to maintain a significant amount of adsorbed hydroxyl species under humid conditions is a key difference as compared to Ni/YSZ electrodes and can potentially affect their electrocatalytic properties. The effect of the gas atmosphere to the electrical conductivity In order to compare the electrical properties of the two electrodes in various gas atmospheres, their ohmic resistance (Rohm) was measured by AC-impedance spectroscopy72. The impedance spectrum contains information about the ionic resistance of the electrolyte as well as the


26

Page 27 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


electronic resistances of the electrodes and the external connections. In our case, the electronic resistance of the counter electrode is not included in the measurements since a cell with a reference electrode was used (see inset draw of figure 8). In addition, the ionic resistance of the YSZ electrolyte can be considered constant, because the temperature, and of course the thickness of the electrolyte, were kept identical throughout all our measurements. Therefore, one can directly relate changes of Rohm with the electrical conductivity of the active material (working electrode). Figure 8 compares the reciprocal of Rohm (hereafter conductance, G) as a function of the examined gas environments, for both Ni/YSZ and LSCrF electrodes. In order to omit systematic errors due to uncertainties in the estimation of electrode thickness and area, electrolyte resistance etc., the conductance is normalized to the maximum recorded value under H2 and O2 atmospheres for Ni/YSZ and LSCrF electrodes respectively. The dependence of G on the gas atmosphere for the two electrodes is remarkably different. In particular, for Ni/YSZ it increases to the maximum in the presence of H2, while it drops to zero in O2. This is directly related to the electronic conductivity of nickel particles which are reduced and oxidized in H2 and O2 respectively as shown by NAP-XPS and XRD results. It is interesting that even though NAP-XPS showed that the metallic Ni concentration in 100% H2O is very low (~7%), the electrical conductance of Ni/YSZ still remains comparable to that of a fully metallic electrode (under H2). This suggests that the impact of the surface NiO layer grown in H2O is not detrimental for the electrical conductivity, which drops considerably only upon thick NiO layer formation (e.g. in O2). However, this contradicts the typical behaviour of Ni/YSZ electrodes during electrolysis under 100% H2O (in the absence of H2), which is known to deactivate fast due to electrode oxidation. This could be explained by the fact that the


27


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 28 of 43

experimental temperature and pressure conditions are too mild to form bulk nickel oxide in H2O and therefore the electronic conductivity is maintained. Alternatively, this might also reflect the fact that the NiO surface layer formed in 100% H2O is electrocatalyticaly less active for water dissociation as compared to Ni, which, apart from conductivity loss, is another reason for , SOC performance deterioration.

Figure 8. Plot of the reciprocal of the ohmic resistance (electrical conductance) of LSCrF and Ni/YSZ electrodes measured at 500°C and at open circuit conditions under the indicated gas atmospheres. In contrast to Ni/YSZ, the LSCrF electrode has the highest electrical conductance in O2 which drops by about 80% in H2 and improves again in presence of humidity. The dependence of LSCrF electrical conductivity to the gas ambient can be explained using arguments from the surface analysis presented above. In oxidative conditions electron holes (h• in the Kröger-Vink


28

Page 29 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


notation) are formed to compensate the charge balance from the substitution of La3+ by Sr2+ during doping, giving rise to a p-type conductivity of the material. The h• are localized to the Cr cations increasing their oxidation state from 3+ to 4+ supporting the results of figure 2e. When the gas atmosphere switches to H2, the Cr4+ cation is reduced to a lower valence and a new charge imbalance is introduced with the creation of more oxygen vacancies according to the equation: ´ 2‫ݎܥ‬஼௥ + ܱ௢ → 1ൗ2 ܱଶ + ܸ௢•• + 2‫ݎܥ‬஼௥ ´ refer to the Cr4+ and Cr3+ cations respectively, where in the Kröger-Vink notation ‫ݎܥ‬஼௥ and ‫ݎܥ‬஼௥

Oo is the lattice oxygen, ܸ௢•• the oxygen vacancies and O2 the gas oxygen. The oxygen vacancy formation is accompanied by a reduced number of holes, which act as the main charge carriers of LSCrF, resulting to the decrease of the electrical conductance as shown in figure 8.73 Please note that the electrical conductance returns to the initial value when the gas is switched back to O2, which is in agreement with the reversible surface modifications found for both electrodes (see fig. 4). Although this model is nicely correlated with the modifications of the Cr L-edge spectra shown in figure 2e, surprisingly, the O 1s spectra do not show any evident modifications between O2 and H2 conditions. To our estimation this indicates that oxygen vacancies formation takes place mainly at the bulk of the electrode, which is not visible in the outer 2 nm observed by the O 1s spectra. This can also explain why the modifications of Cr species are observed only in the Cr Ledge which has deeper probing depth (ca. 5 nm) as compared to Cr 2p NAP-XPS. This suggests that the electrical conductance of the LSCrF electrode is mainly influenced by the conditions in the bulk of the electrode, similar to the observations of the Ni/YSZ. One has to specify here that


29


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 30 of 43

the increase of oxygen vacancies in H2 should be rather limited since the XRD results did not show any evident modification of the bulk structure. Overview of the surface characteristics and comparison of the two electrodes A principal finding of this work is the fast response of the electrode surface composition and oxidation state on the reactive gas environment. It is interesting to compare the near-surface composition given by NAP-XPS in figures 4a and 4b, with respect to the bulk as determined by the post-treatment EDXS analysis included in the same figure. At this point, one should recall that the EDXS analysis is based on the X-rays emitted from the volume of the material and has a typical depth of 1-2 micrometers, which is about 1000 times deeper than NAP-XPS. The EDXS composition of Ni/YSZ and LSCrF electrodes, indicated at the right column in figure 4, is in good agreement with the values expected from the nominal composition. With the exception of the final O2 treatment, the Ni/YSZ displayed an evident enrichment of YSZ in the surface with respect to the bulk composition as determined by EDXS. This reflects the well-known shrinkage effect of nickel grains upon reduction74,75, while those of YSZ are practically unaffected (figure 9). The differences of the initial and final XPS intensity ratio might be related to irreversible modifications of the Ni/YSZ cermet. Previous in situ TEM studies76 have shown the creation of porosity into the nickel particles after reoxidation. One can assume that the porous nickel particles at the final oxidation step have higher surface area which increase relatively the Ni 2p signal and consequently enhances the Ni/Zr ratio. In case of the LSCrF sample, the surface composition (NAP-XPS) is notably different than the bulk (EDXS). This suggests that the characteristics revealed by NAP-XPS analysis are highly localized on the surface of the electrode and do not propagate in the perovskite lattice structure, as verified by the XRD results.


30

Page 31 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Surface changes can also help to elucidate the mechanism of nanoparticles formation on the surface of the LSCrF grains indicated at the SEM images (figure 6e). NAP-XPS shows (figure 4b) a maximum surface enrichment of Sr in H2 and humid atmospheres, while the complementary SEM images show particles with smooth surfaces. This suggests that segregated Sr forms layers, or 2-dimenional particles over the perovskite grains, which enhance considerably the Sr 3d signal as compared to the other perovskite elements located underneath (see also the schematic illustration in figure 9). When the gas atmosphere is switched to O2, the segregated Sr layer agglomerates into large particles (as revealed by SEM images) and the perovskite lattice, situated directly below, becomes exposed to the spectroscopic methods. Direct effect of this agglomeration is the dramatic drop of the Sr 3d peak and the increase of La, Fe and Cr NAP-XPS signals. Theoretical simulations of the photoelectron spectra intensities performed by SESSA software (see supporting information fig. S5), confirmed that if segregated Sr particles are not densely packed, then the Sr 3d peak intensity would be similar to that obtained by a homogeneously mixed LSCrF surface. On the other hand, a 2D surface Sr layer over perovskite can reproduce the observed enhancement of Sr atomic ratio found in the experiment performed on the reduced electrode (fig. 4b). It is interesting to compare these findings with previous studies on perovskite surface chemistry in reactive atmospheres and at elevated temperatures. The effect of oxidative (typically air) atmospheres has been extensively studied on perovskites since this type of electrodes are primarily used in ORR reaction (oxygen electrode). Segregation of Sr over La sites

in

oxidative

environments

is

a

general

characteristic,

for

example

for

La0.8Sr0.2Co0.2Fe0.8O3−δ and SrTiO3 perovskite surfaces.71,77 On the other hand, annealing in reducing conditions (i.e. H2) has been proposed as a method for exsolution of B-site metal


31


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 32 of 43

particles on the surface of the perovskites (e.g. Ni particles for La0.4Sr0.4Sc0.9Ni0.1O3-δ perovskites).78 Here we show that Sr surface concentration is enhanced in reducing/humid atmospheres, surprisingly even more than in the oxidative one. By combining chemical and morphological characteristics of the perovskite electrode we propose that the amount of Sr exposed at the surface is related to the morphology of the segregated Srδ+ and in particular to the rearrangement of Sr from 2D layers to 3D particles. Analysis of lanthanum strontium manganite (LSM) thin-film electrodes has shown that the subsurface segregation of Sr under cathodic polarization is reversible and Sr reappears on the surface at anodic polarization.30 It is also confirmed that the interdiffusion of Sr and La cations from the perovskite lattice towards the surface is a general property which can be stimulated not only by the electrochemical potential, but also by the gas phase. The main characteristics of the two electrode types are schematically summarized in figure 9. As illustrated, the Ni/YSZ electrode undergoes significant morphological and oxidation state changes in response to the gas atmosphere, while for LSCrF, the modifications are mild. However, even if the La and Fe cations of the perovskite lattice are quite stable, at least under the tested conditions, Sr is very mobile and dominates the chemistry and the morphology in the nearsurface region. The depletion of the surface from Sr cations directly influences the Cr cation valence state and consequently the conductivity of the electrode. Another important feature is the ability of Ni/YSZ to accumulate a significant population of adsorbed hydroxyl groups, which was not observed for LSCrF. Finally, the dependence of the electronic conductivity on the gas environment is very different for the two electrodes. As expected Ni/YSZ was more conductive in reducing atmospheres, while LSCrF in oxidative ones. These results make evident that this is


32

Page 33 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


related to the oxidation state of nickel in the case of Ni/YSZ and to the mobility of Sr cations out of the crystal lattice for LSCrF.

Figure 9. Schematic illustration of the surface configuration of Ni/YSZ cermet and LSCrF perovskite electrodes under the various gas atmospheres. The relative ohmic resistance of the electrode in each atmosphere is included. CONCLUSIONS The development of new materials for energy conversion and storage applications necessitates deep understanding of the surface transformations taking place in response to the gas phase environments and to the polarization potential. In this study in situ and post-treatment methods are combined with spectra simulations, in order to investigate how the surface composition, oxidation state and morphology of commonly-used SOC electrodes are affected by the gas environment. The electrode surfaces have been shown to exhibit a dynamic behaviour, relative to the modifications of their features observed under reducing and oxidative reactive mediums. In


33


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 34 of 43

addition, correlation of these microscopic features with the electronic conductivity, a crucial macroscopic characteristic of SOC electrodes, was made. The study presents critical evidence and insight to the understanding of water electrolysis process in SOCs by demonstrating crucial differences to the abundance of the adsorbed hydroxyl groups over the two SOC electrode materials studied. Since the electrocatalyst is a key component of solid oxide electrochemical devices, the presented results may contribute to the development of new more robust electrode materials. ASSOCIATED CONTENT Photograph of the NAP-XPS sample holder, description of the sequence of the NAP-XPS experiments, Cr L3,2 -edge AEY-NEXAFS spectra, the area ratio of the photoelectron peaks related to the A and B cation sites of LSCrF, XRD patterns, depth profile NAP-XPS O 1s spectra, and SESSA software simulations. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author * [email protected] ACKNOWLEDGEMENTS The research leading to these results has received funding from the Fuel Cells and Hydrogen 2 Joint Undertaking under the project SElySOs with grant agreement No 671481. This Joint Undertaking receives support from the European Union’s Horizon 2020 research and innovation programme and Greece, Germany, Czech Republic, France, Norway. Finally, we acknowledge


34

Page 35 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


synchrotron SOLEIL for the allocation of the beamtime and SOLEIL beamline scientists for the use of TEMPO beamline. REFERENCES (1)

Atkinson, a; Barnett, S.; Gorte, R. J.; Irvine, J. T. S.; McEvoy, a J.; Mogensen, M.; Singhal, S. C.; Vohs, J. Advanced Anodes for High-Temperature Fuel Cells. Nat. Mater. 2004, 3 (1), 17–27.

(2)

Yang, L.; Cheng, Z.; Liu, M.; Wilson, L. New Insights into Sulfur Poisoning Behavior of Ni-YSZ Anode from LongTerm Operation of Anode-Supported SOFCs. Energy Environ. Sci. 2010, 3 (11), 1804–1809.

(3)

Graves, C.; Ebbesen, S. D.; Jensen, S. H.; Simonsen, S. B.; Mogensen, M. B. Eliminating Degradation in Solid Oxide Electrochemical Cells by Reversible Operation. Nat Mater 2015, 14 (2), 239–244.

(4)

Yoo, S.; Jun, A.; Ju, Y. W.; Odkhuu, D.; Hyodo, J.; Jeong, H. Y.; Park, N.; Shin, J.; Ishihara, T.; Kim, G. Development of Double-Perovskite Compounds as Cathode Materials for Low-Temperature Solid Oxide Fuel Cells. Angew. Chemie Int. Ed. 2014, 53 (48), 13064–13067.

(5)

Papazisi, K. M.; Balomenou, S.; Tsiplakides, D. Synthesis and Characterization of La0.75Sr0.25Cr0.9M0.1O3 Perovskites as Anodes for CO-Fuelled Solid Oxide Fuel Cells. J. Appl. Electrochem. 2010, 40 (10), 1875–1881.

(6)

Yang, X.; Irvine, J. T. S. (La0.75Sr0.25)0.95Mn0.5Cr0.5O3 as the Cathode of Solid Oxide Electrolysis Cells for High Temperature Hydrogen Production from Steam. J. Mater. Chem. 2008, 18 (20), 2349–2354.

(7)

Qi, W.; Ruan, C.; Wu, G.; Zhang, Y.; Wang, Y.; Xie, K.; Wu, Y. Reversibly in-Situ Anchoring Copper Nanocatalyst in Perovskite Titanate Cathode for Direct High-Temperature Steam Electrolysis. Int. J. Hydrogen Energy 2014, 39 (11), 5485–5496.

(8)

Neofytides, C.; Athanasiou, M.; Neophytides, S. G.; Niakolas, D. K. Sulfur Tolerance of Au-Mo-Ni/GDC SOFC Anodes under Various CH4 Internal Steam Reforming Conditions. ECS Trans. 2015, 68 (1), 1353–1360.

(9)

Ruiz-Morales, J. C.; Marrero-López, D.; Canales-Vázquez, J.; Irvine, J. T. S. Symmetric and Reversible Solid Oxide Fuel Cells. RSC Adv. 2011, 1 (8), 1403–1414.

(10)

Myung, J. H.; Neagu, D.; Miller, D. N.; Irvine, J. T. Switching on Electrocatalytic Activity in Solid Oxide Cells. Nature 2016, 537 (7621), 528–531.


35


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(11)

Page 36 of 43

Ding, H.; Tao, Z.; Liu, S.; Yang, Y. A Redox-Stable Direct-Methane Solid Oxide Fuel Cell (SOFC) with Sr2FeNb0.2Mo0.8O6????? Double Perovskite as Anode Material. J. Power Sources 2016, 327, 573–579.

(12)

M. Kubota, H. Muroyama, T. Matsui, and K. E. Microstructural Change of Ni–YSZ Anode under Thermal Cycles with Redox Treatments. Electrochem. Soc. 2013, 53 (1), 1689–1699.

(13)

Ettler, M.; Timmermann, H.; Malzbender, J.; Weber, A.; Menzler, N. H. Durability of Ni Anodes during Reoxidation Cycles. J. Power Sources 2010, 195 (17), 5452–5467.

(14)

Feng, Z.; Hong, W. T.; Fong, D. D.; Lee, Y.-L.; Yacoby, Y.; Morgan, D.; Shao-Horn, Y. Catalytic Activity and Stability of Oxides: The Role of Near-Surface Atomic Structures and Compositions. Acc. Chem. Res. 2016, 49 (5), 966–973.

(15)

Liu, X.; Yang, W.; Liu, Z. Recent Progress on Synchrotron-Based In Situ Soft X-Ray Spectroscopy for Energy Materials. Adv. Mater. 2014, 1–20.

(16)

El Gabaly, F.; McCarty, K. F.; Bluhm, H.; McDaniel, A. H. Oxidation Stages of Ni Electrodes in Solid Oxide Fuel Cell Environments. Phys. Chem. Chem. Phys. 2013, 15 (21), 8334–8341.

(17)

Papaefthimiou, V.; Niakolas, D. K.; Paloukis, F.; Dintzer, T.; Zafeiratos, S. Is Steam an Oxidant or a Reductant for Nickel/Doped-Ceria Cermets? ChemPhysChem 2017, 18 (1), 164–170.

(18)

Law, Y. T.; Zafeiratos, S.; Neophytides, S. G.; Orfanidi, A.; Costa, D.; Dintzer, T.; Arrigo, R.; Knop-Gericke, A.; Schlögl, R.; Savinova, E. R. In Situ Investigation of Dissociation and Migration Phenomena at the Pt/electrolyte Interface of an Electrochemical Cell. Chem. Sci. 2015, 6 (10), 5635–5642.

(19)

Tsvetkov, N.; Lu, Q.; Sun, L.; Crumlin, E. J.; Yildiz, B. Improved Chemical and Electrochemical Stability of Perovskite Oxides with Less Reducible Cations at the Surface. Nat. Mater. 2016, 15 (9), 1010–1016.

(20)

Feng, Z.; Crumlin, E. J.; Hong, W. T.; Lee, D.; Mutoro, E.; Biegalski, M. D.; Zhou, H.; Bluhm, H.; Christen, H. M.; Shao-Horn, Y. In Situ Studies of the Temperature-Dependent Surface Structure and Chemistry of Single-Crystalline (001)-Oriented La0.8Sr0.2CoO3-δ Perovskite Thin Films. J. Phys. Chem. Lett. 2013, 4 (9), 1512–1518.

(21)

Bozzini, B.; Amati, M.; Bocchetta, P.; Dal Zilio, S.; Knop-Gericke, A.; Vesselli, E.; Kiskinova, M. An in Situ nearAmbient Pressure X-Ray Photoelectron Spectroscopy Study of Mn Polarised Anodically in a Cell with Solid Oxide Electrolyte. Electrochim. Acta 2015, 174, 532–541.

(22)

Opitz, A. K.; Nenning, A.; Rameshan, C.; Rameshan, R.; Blume, R.; Haevecker, M.; Knop-Gericke, A.; Rupprechter,


36

Page 37 of 43

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


G.; Fleig, J.; Kloetzer, B. Enhancing Electrochemical Water-Splitting Kinetics by Polarization-Driven Formation of Near-Surface Iron(0): An In Situ XPS Study on Perovskite-Type Electrodes. Angew. Chemie-International Ed. 2015, 54 (9), 2628–2632.

(23)

Zhou, J.; Shin, T. H.; Ni, C.; Chen, G.; Wu, K.; Cheng, Y.; Irvine, J. T. S. In Situ Growth of Nanoparticles in Layered Perovskite La0.8Sr1.2Fe0.9Co0.1O4-δ as an Active and Stable Electrode for Symmetrical Solid Oxide Fuel Cells. Chem. Mater. 2016, 28 (9), 2981–2993.

(24)

Papaefthimiou, V.; Shishkin, M.; Niakolas, D. K.; Athanasiou, M.; Law, Y. T.; Arrigo, R.; Teschner, D.; Hävecker, M.; Knop-Gericke, A.; Schlögl, R.; Ziegler, T.; Neophytides, S. G.; Zafeiratos, S. On the Active Surface State of NickelCeria Solid Oxide Fuel Cell Anodes during Methane Electrooxidation. Adv. Energy Mater. 2013, 3 (6), 762–769.

(25)

Crumlin, E. J.; Liu, Z.; Bluhm, H.; Yang, W.; Guo, J.; Hussain, Z. X-Ray Spectroscopy of Energy Materials under in Situ/operando Conditions. J. Electron Spectros. Relat. Phenomena 2015, 200, 264–273.

(26)

Opitz, A. K.; Nenning, A.; Rameshan, C.; Rameshan, R.; Blume, R.; Hävecker, M.; Knop-Gericke, A.; Rupprechter, G.; Fleig, J.; Klötzer, B. Enhancing Electrochemical Water-Splitting Kinetics by Polarization-Driven Formation of nearSurface iron(0): An in Situ XPS Study on Perovskite-Type Electrodes. Angew. Chemie - Int. Ed. 2015, 54 (9), 2628– 2632.

(27)

Nenning, A.; Opitz, A. K.; Rameshan, C.; Rameshan, R.; Blume, R.; Hävecker, M.; Knop-Gericke, A.; Rupprechter, G.; Klötzer, B.; Fleig, J. Ambient Pressure XPS Study of Mixed Conducting Perovskite-Type SOFC Cathode and Anode Materials under Well-Defined Electrochemical Polarization. J. Phys. Chem. C 2016, 120 (3), 1461–1471.

(28)

Hong, W. T.; Stoerzinger, K. A.; Crumlin, E. J.; Mutoro, E.; Jeen, H.; Lee, H. N.; Shao-Horn, Y. Near-Ambient Pressure XPS of High-Temperature Surface Chemistry in Sr2Co2O5 Thin Films. Top. Catal. 2016, 59 (5–7), 574–582.

(29)

Crumlin, E. J.; Mutoro, E.; Liu, Z.; Grass, M. E.; Biegalski, M. D.; Lee, Y.-L.; Morgan, D.; Christen, H. M.; Bluhm, H.; Shao-Horn, Y. Surface Strontium Enrichment on Highly Active Perovskites for Oxygen Electrocatalysis in Solid Oxide Fuel Cells. Energy Environ. Sci. 2012, 5 (3), 6081–6088.

(30)

Huber, A.-K.; Falk, M.; Rohnke, M.; Luerssen, B.; Amati, M.; Gregoratti, L.; Hesse, D.; Janek, J. In Situ Study of Activation and de-Activation of LSM Fuel Cell Cathodes – Electrochemistry and Surface Analysis of Thin-Film Electrodes. J. Catal. 2012, 294, 79–88.

(31)

Blume, R.; Hävecker, M.; Zafeiratos, S.; Teschner, D.; Kleimenov, E.; Knop-Gericke, A.; Schlögl, R.; Barinov, A.;


37


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 38 of 43

Dudin, P.; Kiskinova, M. Catalytically Active States of Ru(0001) Catalyst in CO Oxidation Reaction. J. Catal. 2006, 239 (2), 354–361.

(32)

Lu, Q.; Chen, Y.; Bluhm, H.; Yildiz, B. Electronic Structure Evolution of SrCoOx during Electrochemically Driven Phase Transition Probed by in Situ X-Ray Spectroscopy. J. Phys. Chem. C 2016, 120 (42), 24148–24157.

(33)

Papaefthimiou, V.; Niakolas, D. K.; Paloukis, F.; Teschner, D.; Knop-Gericke, A.; Haevecker, M.; Zafeiratos, S. Operando Observation of Nickel/ceria Electrode Surfaces during Intermediate Temperature Steam Electrolysis. J. Catal. 2017, 352, 305–313.

(34)

Zhang, C.; Grass, M. E.; McDaniel, A. H.; DeCaluwe, S. C.; El Gabaly, F.; Liu, Z.; McCarty, K. F.; Farrow, R. L.; Linne, M. A.; Hussain, Z.; Jackson, G. S.; Bluhm, H.; Eichhorn, B. W. Measuring Fundamental Properties in Operating Solid Oxide Electrochemical Cells by Using in Situ X-Ray Photoelectron Spectroscopy. Nat. Mater. 2010, 9 (11), 944– 949.

(35)

Huber, A. K.; Falk, M.; Rohnke, M.; Luerssen, B.; Gregoratti, L.; Amati, M.; Janek, J. In Situ Study of Electrochemical Activation and Surface Segregation of the SOFC Electrode Material La0.75Sr0.25Cr0.5Mn0.5O3±δ. Phys Chem Chem Phys 2012, 14 (2), 751–758.

(36)

Bastidas, D. M.; Tao, S.; Irvine, J. T. S. A Symmetrical Solid Oxide Fuel Cell Demonstrating Redox Stable Perovskite Electrodes. J. Mater. Chem. 2006, 16 (17), 1603–1605.

(37)

Sapountzi, F. M.; Brosda, S.; Papazisi, K. M.; Balomenou, S. P.; Tsiplakides, D. Electrochemical Performance of La0.75Sr0.25Cr0.9M0.1O3 Perovskites as SOFC Anodes in CO/CO2 Mixtures. J. Appl. Electrochem. 2012, 42 (9), 727–735.

(38)

Hansen, K. K.; Menon, M.; Knudsen, J.; Bonanos, N.; Mogensen, M.; Kammer Hansen, K.; Menon, M.; Knudsen, J.; Bonanos, N.; Mogensen, M. The Effect of a CGO Barrier Layer on the Performance of LSM/YSZ SOFC Cathodes. J. Electrochem. Soc. 2010, 157 (3), B309–B313.

(39)

Yeh, J. J.; Lindau, I. Atomic Subshell Photoionization Cross Sections and Asymmetry Parameters: 1

Insights into the Surface Reactivity of Cermet and Perovskite

Recommend Documents