Enhanced oxygen mobility by doping Yb in BaGd1-xYbxMn2O5+δ double perovskite-structured oxygen storage materials

Solid State Ionics 335 (2019) 103–112

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Enhanced oxygen mobility by doping Yb in BaGd1-xYbxMn2O5+δ double perovskite-structured oxygen storage materials Kun Zheng

T

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AGH University of Science and Technology, Faculty of Energy and Fuels, Department of Hydrogen Energy, al. A. Mickiewicza 30, 30-059 Krakow, Poland AGH University of Science and Technology, AGH Centre of Energy, ul. Czarnowiejska 36, 30-054 Krakow, Poland

A R T I C LE I N FO

A B S T R A C T

Keywords: Oxygen storage materials A-site ordered double perovskite BaGd1-xYbxMn2O5+δ Oxygen nonstoichiometry and transport properties Oxygen diffusion coefficient

The A-site ordered BaGd1-xYbxMn2O5+δ double perovskites have been systematically investigated in terms of the crystal structure, oxygen nonstoichiometry, microstructure and oxidation states of Mn, oxygen storage properties and oxygen in-situ intake, as well as transport properties. The increased content of Yb in BaGd1-xYbxMn2O5+δ results in a decrease of the relative unit cell volume changes between oxidized materials and reduced counterparts. Crystal structure with P4/nmm space group is observed at room temperature in all BaGd1-xYbxMn2O5+δ (x = 0, 0.2 and 0.4), except in BaGd0.6Yb0.4Mn2O6. The increased content of Yb in BaGd1-xYbxMn2O5+δ modifies materials' morphology, allowing to obtain fine powders. The Mn2+ and/or Mn3+ oxidation states are present in the reduced BaGd1-xYbxMn2O5, confirmed by XPS studies. The Yb doping in BaGd0.6Yb0.4Mn2O5+δ significantly decreases the reduction time and oxidation/reduction temperature. The in-situ oxidation of BaGd0.6Yb0.4Mn2O5 in air associated with P4/nmm → P-1 space group change occurs between 225 and 275 °C, and its nature may indicate the oxygen diffusion mainly occurs in the Gd0.6Yb0.4 layers. The TEC values calculated from high temperature XRD data, not exceeding 14.8(1) × 10−6 K−1, are moderate. The high electrical conductivity (100 S cm−1 in air at 600 °C) and high oxygen diffusion coefficient suggest wide potential applications of the developed BaGd0.6Yb0.4Mn2O5+δ.

1. Introduction The oxygen storage materials (OSMs) in the past years have attracted a lot of great scientific research interest, and this is due to the possibility of the commercial application in various technologies and industrial processes, for instance: solid oxide fuel cell (SOFC) electrodes, automotive applications (exhaust gas treatment), oxygen selective membranes, thermochemical CO2 and/or H2O splitting for fuel production, non-aerobic oxidation including flameless combustion of hydrocarbons, high-temperature production that requires high-purity oxygen, oxy-fuel and chemical looping combustion processes used in clean coal-type energy production, production of synthesis gas, inert gas purification due to oxygen scavenger behavior, etc. [1–9]. The availability of stable oxygen storage materials with a large oxygen storage capacity (OSC), fast and reversible oxygen intake/release kinetic at low working temperature is very crucial for the application of the mentioned technologies. Until now, the most widely applied oxygen storage materials are the Ce1-xZrxO2-δ–based compounds with the fluorite structure and oxygen storage capacities are of about 1500 μmolO/g [10,11]. Extremely high oxygen storage capacity with 2700 μmol-

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O/g has been detected in hexagonal-structured YBaCo4O7+δ oxide, while the Co-based material suffers from severe stability issues [12], and YBaCo4O7+δ-based modified oxides have also been proposed [13]. Other materials, such as: layered structured LuFe2O4+δ materials with OSC of ~1400 μmol-O/g [14], hexagonal structured Dy1-xYxMnO3+δ with OSC of ~1200 μmol-O/g [15,16], Ca2AlMnO5+δ with brownmillerite structure and OSC of ~1900 μmol-O/g [17], Ln2O2S – Ln2O2SO4 (Ln: La, Pr, Nd and Sm) systems [18,19] and simple perovskites (such as: SrFeO3-δ and Cu-containing derivatives, (Sr,Ca)MnO3-δ oxides) [6], have also been investigated. One of the most promising novel oxygen storage materials with very high OSC is the series of BaLnMn2O5+δ (Ln: Pr, Nd, Sm, Gd, Dy, Er and Y; OSC of ~2400 μmol-O/g) double perovskites. This system shows reduced temperatures of reduction and oxidation processes, and fast kinetics of reversible oxygen content δ changes occurring at low temperatures (300 °C–500 °C) with a change of the oxygen partial pressures (e.g. air and 5 vol% H2 in argon) [3,8,9,20–22]. The structure of BaLnMn2O5+δ can be described as an Asite ordered double perovskite with layers of Ln and barium ions in the A-site alternating along [001] [3,8,9,20,21]. The unit cell in this case is created by a doubling of the simple perovskite lattice constant ap along

AGH University of Science and Technology, Faculty of Energy and Fuels, Department of Hydrogen Energy, al. A. Mickiewicza 30, 30-059 Krakow, Poland. E-mail address: [email protected].

https://doi.org/10.1016/j.ssi.2019.03.001 Received 26 November 2018; Received in revised form 15 February 2019; Accepted 1 March 2019 0167-2738/ © 2019 Elsevier B.V. All rights reserved.

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ppm (10−12) using ZIROX SGM5 oxygen analyzer working at 750 °C. The used synthesis conditions have been carefully chosen for the synthesis on a basis of previous experiments with BaLnMn2O5+δ (Ln – lanthanides, Y) system [8,9,20]. Synthesis of BaGd1-xYbxMn2O5+δ (x = 0.5–1) samples was also attempted, however, a multiphase product was always obtained, with presence of BaMnO3 and/or YbMnO3 phases, among others. This is a crucial result, showing the doping limit of Yb in BaGd1-xYbxMn2O5+δ is restricted to x = 0.4, and there is a strong tendency of formation of BaMnO3 and YbMnO3. In the case of BaGd0.8Yb0.2Mn2O5+δ and BaGd0.6Yb0.4Mn2O5+δ compounds, a very small amount of additional BaMnO3-like impurities was detected. Structural studies at room temperature of the synthesized materials were performed in the 10–110° range with CuKα radiation, using Panalytical Empyrean diffractometer. The oxygen nonstoichiometry of oxidized BaGd1-xYbxMn2O5+δ materials has also been determined by iodometric titration of Mn using EasyPlus Titrator equipment. For the measurements, about 0.15 g of the studied compounds was dissolved in 1 M HCl solution, which contains 1 g KI (excess amount). The Mn3+ and Mn4+ cations in materials are reduced to Mn2+, and as a result a stoichiometric amount of iodine was formed. The prepared solution was then titrated with 0.1 M Na2S2O3 solution, which was previously standardized with K2Cr2O7. The accruing process can be descried by the following reactions [33]:

c-axis, which is often noted as 1ap × 1ap × 2ap, or simply as 1 × 1 × 2 [23]. The most simple ordering occurs as 1:1-type, and can be present in A sublattice (AA'B2O6) or/and B-sublattice (A2BB'O6, AA'BB'O6) [23,24]. Crystal structure of such cation-ordered compounds originates from the simple perovskite one, and can be constructed by an appropriate consideration of layer-, columnar- or rock salt-type cation ordering [23]. In the series of BaLnMn2O6 oxygen storage materials, the reduction is the limiting factor (as the oxidation proceeds much faster) of oxygen storage-related performance. From the practical perspective of applications, it is crucial to develop materials with enhanced reduction kinetics [3,8,9,22]. Interestingly, the fastest kinetics in the series of BaLnMn2O5+δ double perovskites were measured during reduction of BaGdMn2O6, which indicates the modification of BaGdMn2O5+δ-based material is of great interest to improve the oxygen storage properties. Furthermore, the unit cell volume for series of reduced BaLnMn2O5 and oxidized BaLnMn2O6 perovskites was found to be linearly dependent (with different slope) on the ionic radius of Ln3+ cation, respectively. The smallest relative change of unit cell volume between BaLnMn2O5 and BaLnMn2O6, is recorded in the BaErMn2O5+δ with the smallest Er3+ in the series of Ln3+ [8,9]. However, the synthesis attempts of BaYbMn2O5+δ with even smaller ionic radius of Yb3+ were not successful [8,9,25]. The chemical composition modification of A-site cations in BaLnMn2O5+δ series improved the oxygen storage related properties, for instance: Ba0.9Sr0.1YMn2O5+δ material [26], BaY1xGdxMn2O5+δ oxides [27], BaY1-xPrxMn2O5+δ compounds [28], BaY1xSmxMn2O5+δ perovskites [26,29]. However, no clear tendency between size or electronegativity of the substituted cation and reduction kinetic (for the same morphology of investigated samples) was reported. Other approaches, such as: high energy milling [30,31] and wet-chemical synthesis method [32] were also applied to enhance the oxygen intake/release kinetics by morphology modifications of the oxygen storage materials powder. In this work, substitutions of Gd with small cation Yb3+ at A-site in BaGdMn2O5+δ parent material have been performed, and the reuslts show the oxygen mobility is enhanced by doping Yb in BaGd1xYbxMn2O5+δ double perovskite materials. The crystal structure, oxygen nonstoichiometry, materials' microstructure, oxygen storagerelated properties, the in-situ oxygen intake investigated by high temperature XRD measurements, and the transport properties have been systematically studied.

Reduction:2Mn3 + + 2I− → 2Mn2 + + I2 ; Mn4 + + 2I− → Mn2 + + I2

Titration:I2 + 2S2 O32 − → 2I− + S4 O62 − The whole titration process was performed in argon atmosphere, and a EM40-BNC Mettler Toledo combined platinum electrode with Ag/ AgCl reference element was used to determine the finish point of the titration. The Scanning Electron Microscopy (SEM) measurements were performed using FEI Nova NanoSEM 200 apparatus on powders obtained after crushing and grinding of “as synthesized” pellets. Additional particle size analysis of the powders of BaGd1-xYbxMn2O5+δ (x = 0, 0.2 and 0.4) has been performed by Mastersizer 3000 laser diffraction particle size analyzer. XPS (X-ray photoelectron spectroscopy) spectra of BaGd1-xYbxMn2O5 materials were acquired on an Axis UltraDLD spectrometer (Kratos Analytical) using monochromatic Al Kα radiation. Oxygen storage-related measurements, including reduction/oxidation runs, were conducted by the thermogravimetric (TG) method. All experiments were performed on TA Q5000IR apparatus. Measurements were done on powdered samples, obtained after grinding of sinters and sieving on 100 μm sieve. The reduction process was carried out in the flow of 5 vol% H2 in Ar, while the synthetic air was used for the oxidation process. Isothermal reduction/oxidation runs were recorded at 500 °C and 600 °C, respectively, and the data shown in graphs concern 5th cycle. Non-isothermal oxidations and reductions were collected after these five isothermal cycles. Calculation of the reversible oxygen storage capacity was corrected with buoyancy effect, which was established on a basis of runs performed without powder samples (TG pan only). The in-situ studies of oxygen intake of BaGd1-xYbxMn2O5 samples was investigated by high temperature XRD measurements in synthetic air flow, in 30–800 °C temperature range. While the in-situ investigations of oxygen release process couldn't be performed due to the technical limitation of XRD equipment. High temperature XRD experiments were done on Panalytical Empyrean apparatus equipped with Anton Paar HTK 1200 N oven-chamber. Rietveld analysis of all XRD data was done in GSAS/EXPGUI set of software [34]. Minimal offset of data gathered using room temperature holder and high temperature holder (measurements at 30 °C) was small enough to be neglected. Results of high temperature XRD data were also used for calculation of thermal expansion coefficient (TEC) of the considered materials. A linear TEC behavior was derived with an isotropic expansion approximation, as 1/ 3 of the volumetric thermal expansion coefficient, for which refined

2. Experimental In order to avoid the misperception of naming scheme in this work, the following convention is used: for the reduced samples, oxygen content is marked as being close to 5 (BaGd1-xYbxMn2O5+δ, δ ≈ 0), while in the case of oxidized materials it is labelled as BaGd1xYbxMn2O6 (δ ≈ 1). In the discussion of oxygen storage-related properties the chemical composition of the material is written as BaGd1xYbxMn2O5+δ. BaGd1-xYbxMn2O5+δ (x = 0, 0.2 and 0.4) compounds were successfully synthesized by soft chemistry method. After the dissolution of stoichiometric amounts of respective nitrates in deionized water, the ammonia salt of ethylenediaminetetra-acetic acid (used as the complexing agent) was added into the solutions. The obtained homogeneous solutions were heated in quartz containers up to around 400 °C. The evaporation of water, decomposition of excessive ammonia nitrate and oxidation of residual carbon occurred during the heating. Final heat treatment of the materials was conducted at 1100 °C for 8 h in 1 vol% hydrogen in argon. The applied 1 vol% H2 in argon in the synthesis was prepared by mixing of 5 vol% H2 in Ar (AirLiquid) with 5 N Ar gas (AirLiquid, O2 content ≤2.0 ppm-mol) with respective flows set on mass flow controllers. The mixing was done to obtain ca. 1 vol% hydrogen in argon mixture, and the oxygen partial pressure in the mixture was measured (at the outlet of the furnace) to be ca. 10−6 vol104

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used to refine the XRD data with achieving good refinement results. While for the oxidized BaGd0.6Yb0.4Mn2O6 (see Fig. 1b) material, the actual symmetry is much lower. In this case, the triclinic P-1 space group is applied. It is reported that in the series of BaLnMn2O6 double perovskites, space groups with low symmetry (monoclinic P121 or triclinic P-1) are used to refine materials with small Ln3+ (Dy3+, Y3+ and Er3+) cations [8,9,22]. The average ionic radius of Gd0.6Yb0.4 in BaGd0.6Yb0.4Mn2O5+δ is very close to that of Dy in BaDyMn2O6, as show in Fig. 2b. In this work, the choice of P-1 space group (BaDyMn2O6 shows P-1 structure [8]) for the refinement of oxidized BaGd0.6Yb0.4Mn2O6 oxide is reasonable. By normalizing unit cell parameters to ap, c/a ratio for BaGd0.6Yb0.4Mn2O5 is equal to 0.9745, while for BaGd0.6Yb0.4Mn2O6 material c/a equals 0.9743 and c/b = 0.9749. As all angles of P-1unit cell are very close to 90 deg., such a calculation may preliminarily suggest a rather uniform shrinkage/elongation of the material upon oxidation/reduction. In order to elucidate the oxidation process, in situ studies of the BaGd0.6Yb0.4Mn2O5 oxidation process by high temperature XRD are performed and the results are discussed in the following section. The continuous doping of Yb (x = 0.5) leads to a large presence of impurities, as shown in Fig. 2a, and even several times of synthesis repeat in the same condition with intermediate groundings cannot decrease the amount of impurities. Moreover, the synthesis of BaYbMn2O5+δ has also been tried and the result showed the material is totally decomposed to BaMnO3 and YbMnO3 phases. This indicates that the doing limit of Yb in BaGd1-xYbxMn2O5+δ is restricted to x = 0.4, and a strong tendency of the formation of BaMnO3 and YbMnO3. The respective unit cell volume data for the reduced and oxidized BaGd1xYbxMn2O5+δ materials follow the observed trend of changes in the whole BaLnMn2O5+δ series closely (see Fig. 2b). The rest BaLnMn2O5+δ data used in Fig. 2b are from the author's previous co-work [9]. The relative change of unit cell volumes of BaGd1-xYbxMn2O5+δ, calculated as a difference between values for reduced and oxidized materials and divided by values for the reduced counterparts, respectively, is also shown in the Fig. 2b. Although the relative unit cell volume change of BaGd0.6Yb0.4Mn2O5+δ is not the smallest, the further doping of Yb cannot be successful. This may be related with the fact that the very small cation Yb has a tendency to form YbMnO3 phase instead of occupying the Gd-site in double perovskite structure. The relative unit cell volume change between oxidized BaGd0.6Yb0.4Mn2O6 and reduced BaGd0.6Yb0.4Mn2O5 is 1.9%, and it is worth noting that such value is much lower than the NiO to Ni volume contraction of ~40% in the typical anode material for SOFCs [39]. This suggests that the relative volume change in BaGd1-xYbxMn2O5+δ compounds will not be a problem for the wide potential application of BaGd1-xYbxMn2O5+δ materials in reducing and oxidizing conditions, such as electrode materials for SOFCs [1,9,39,40].

values of the unit cell volume of the samples were taken. The electrical conductivity changes of BaGd0.6Yb0.4Mn2O5+δ sample were also recorded during the oxygen release/intake process (reduction/oxidation process). The conductivity measurements were performed in 5 vol% H2/Ar and air in 30–850 °C temperature range on a custom-made equipment. The conductivity changes were also recorded during the reduction and oxidation process. For the studies, dense sinters were cut into cuboid shape with approximate dimensions of about 3 × 3 × 10 mm. A four-probe DC method was applied for determination of total conductivity σ. For the external electrodes Au plates were used, while the internal ones made from thin Pt wire, were wrapped several times around the sinter. The oxygen diffusion coefficient D and surface exchange coefficient k of BaGd0.6Yb0.4Mn2O6 oxide were also evaluated by the mass relaxation technique [35,36]. The measurements were performed on TA Instruments Q5000 IR apparatus on thin-sheet shape sinter. The mass relaxation related to change of the oxygen stoichiometry in the sample was recorded during rapid change (flush time < 1 s) of oxygen partial pressure between 0.1 atm and 0.01 atm at 500 °C. Calculation of D and k was done in custom-made Matlab code, based on 1-dimensional solution for thin-sheet shaped sample provided by Crank [35,37].

3. Results and discussion 3.1. Crystal structure of BaGd1-xYbxMn2O5+δ The crystal structural data obtained by the refinement for reduced BaGd1-xYbxMn2O5 and oxidized BaGd1-xYbxMn2O6 are gathered in Table 1. The recorded diffractograms of reduced BaGdMn2O5 and oxidized BaGdMn2O6 compounds can be successfully refined using tetragonal P4/nmm aristotype space group, as reported in the literature [8,21,22,38]. The obtained unit cell parameters and cell volume of BaGdMn2O5 and BaGdMn2O6 are very similar to the data published previously [8]. Interestingly, a small amount of Yb doping (x = 0.2) doesn't change the crystal structure of BaGd0.8Yb0.2Mn2O5 (see Fig. 1a) and BaGd0.8Yb0.2Mn2O6 materials, and the good refinement of the XRD data can be obtained still assuming P4/nmm (√2ap × √2ap × 2ap) superstructure. However, a small amount of BaMnO3 secondary phase (about 1.2 wt%) was detected in BaGd0.8Yb0.2Mn2O5+δ material. In the P4/nmm structure of BaGd1-xYbxMn2O5+δ (x = 0, 0.2), the Mn cations at B-site form rock salt-like charge order with (formal) Mn2+ and Mn3+ (or Mn3+ and Mn4+ for oxidized materials) states at different crystallographic positions [8,23]. The continuous doping of Yb (x = 0.4) into A-site of BaGd1-xYbxMn2O5+δ changes the crystal structure, and about 4.5 wt% BaMnO3 secondary phase is also present (the amount of YbMnO3 is negligible). In the case of reduced BaGd0.6Yb0.4Mn2O5 (see Fig. 1b) compound, the tetragonal P4/nmm space group can still be Table 1 Rietveld refinement results for BaGd1-xYbxMn2O5+δ (x = 0, 0.2, 0.4) oxides. Composition

Space group a [Å]

1

BaGdMn2O5+δ

2

BaGd0.8Yb0.2Mn2O5+δ

Oxidized (O5.92)

Reduced (O5)

Oxidized (O5.97)

Reduced (O5)

Oxidized (O5.93)

Reduced (O5)

P4/nmm 5.5319(1)

P4/nmm 5.5812(1)

P4/nmm 5.5347 (1)

P4/nmm 5.5736(1)

P4/nmm 5.5605(1)

7.6294(1)

7.6880(1)

7.6123(1)

7.6777(1)

233.47(1) 1.20 3.39 4.43

239.48(1) 1.38 2.92 3.92

233.19(1) 1.52 3.2 4.27

238.51(1) 1.37 3.02 4.03

P-1 5.5291(1) 90.01 deg 5.5274(1) 90.27 deg 7.6207(1) 89.98 deg 232.63(1) 1.24 2.86 3.71

b [Å] c [Å] 3

V [Å ] CHI2 Rp (%) Rwp (%) 1, 2

BaGd0.6Yb0.4Mn2O5+δ

contain 1.2 wt% and 4.5 wt% of BaMnO3-like impurities, respectively. 105

7.6696(1) 237.14(1) 1.76 3.63 4.98

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Fig. 1. XRD diffractogram with Rietveld analysis for a) BaGd0.8Yb0.2Mn2O5, b) BaGd0.6Yb0.4Mn2O5 and BaGd0.6Yb0.4Mn2O6 materials (with a very small amount of BaMnO3-like impurities).

Fig. 2. a) Synthesis of BaGd0.5Yb0.5Mn2O5+σ and BaYbMn2O5+σ, b) normalized unit cell volume of BaGd1-xYbxMn2O5 and BaGd1-xYbxMn2O6 as a function of Ln ionic radius (BaLnMn2O5+σ data are from the author's previous work [9]).

As shown in the Table 1, the increased content of small Yb3+ cation in BaGd1-xYbxMn2O5+δ compounds decreases the unit cell volumes. The reduced materials possess much larger unite cell volumes, as compared to the oxidized ones, respectively. This behavior can be related with several factors: a). an increase of average radius of Mn cations after reduction, associated with a decrease of average oxidation states from +3.5 to +2.5; b). an increase of ionic radius with a decrease of coordination number, which is equal to 5 in square pyramidal coordination of manganese in BaGd1-xYbxMn2O5 and 6 in octahedral coordination of Mn in BaGd1-xYbxMn2O6; c). weaker bonding in BaGd1xYbxMn2O5 oxides due to the oxygen empty layer of Gd1-xYbx. The increased unit cell volumes can also be related with the influence of chemical expansion, usually observed in simple ABO3-δ perovskites, and the partially reduced samples exhibit larger unit cell volumes with the increasing oxygen vacancies [41]. The simple geometric shrinkage due to the lack of oxygen anions in BaGd1-xYbxMn2O5 compounds is not dominant. As can be seen in Fig. 2b, the reduced BaLnMn2O5 materials (also BaGd1-xYbxMn2O5) possess much larger unit cell volume in

relation to the corresponding oxidized ones. In this case the simple geometric shrinkage due to the lack of oxygen anions cannot be dominant. This can be simply understood as the chemical expansion, which is very often observed at high temperatures for Co-containing perovskites [41]. The layered-type cation ordering of BaGd1-xYbx arrangement in BaGd1-xYbxMn2O5+δ originates mainly from the ionic size difference between Ba2+ and Gd/Yb3+ cations, with charge difference not being very crucial [23]. While one has to note that for the serials of BaLnMn2O5+δ oxygen storage materials, the cation-disordering or cation-ordering for the same chemical composition is strongly related with the synthesis conditions, if the Ln radius is not very small [1,38]. Crystal structure data for all reduced BaGd1-xYbxMn2O5 (x = 0, 0.2 and 0.4) compounds were refined using P4/nmm symmetry, indicating Bsite rock salt-like charge order of Mn2+ and Mn3+ cations. While there is no electronic phase diagram available for all BaGd1-xYbxMn2O5 compounds, supposedly all these materials possess significantly lower electrical conductivity than their oxidized ones (see the conductivity data), suggesting localized charges and supporting the choice of P4/ 106

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Fig. 3. Scanning electron micrographs of (a) BaGdMn2O5, (b) BaGd0.8Yb0.2Mn2O5, (c) BaGd0.6Yb0.4Mn2O5 materials; and particle size analysis results of (d) BaGdMn2O5, (e) BaGd0.8Yb0.2Mn2O5, (f) BaGd0.6Yb0.4Mn2O5 materials.

nmm space group with different sites for Mn2+ and Mn3+ cations. Similar crystal structure data were also observed for BaLnMn2O5 oxides in the literature [8,9,21]. The oxygen content of oxidized BaGd1-xYbxMn2O5+δ materials has been determined by iodometric titration of Mn. The oxygen non-stoichiometry in the oxidized BaGd1-xYbxMn2O5+δ materials was determined as: 5.92(1), 5.97(1) and 5.93(1), respectively (for x = 0, 0.2 and 0.4 samples). Moreover, the non-stoichiometry of fully reduced samples was also calculated based on the TG data, and the oxygen content was found to be close to 5.0 (slightly below). Overall, the evaluated oxygen content and its changes of the studied BaGd1xYbxMn2O5+δ materials are similar to the reported values for BaLnMn2O5+δ and similar doped materials [21,42,43].

increase of small particles amount and a decrease of the large particles. In the BaGd0.6Yb0.4Mn2O5 compound, the presence of large particles with 100 μm disappears, and the small particles are dominant in the powder, which favor the oxygen storage performance. An attempt was also made to investigate the oxidation states of Mn cations in the considered reduced BaGd1-xYbxMn2O5 (x = 0, 0.2 and 0.4) materials by XPS studies. The measured XPS data of Mn 2p can be well fitted with MnO/Mn2O3 references (Fig. 4). The binding energy of MnO and Mn2O3 are overlapped due to the similarity of signals from manganese cations at oxidation states of Mn2+ and Mn3+ [45]. As can be seen in Fig. 4, the deconvoluted peaks at 641.7 and 653.4 eV (for BaGdMn2O5 sample), 641.5 and 653.4 eV (for BaGd0.8Yb0.2Mn2O5), and 641.7 and 653.4 eV (for BaGd0.6Yb0.4Mn2O5 sample) belong to Mn2+/

3.2. Microstructure studies and XPS results The microstructures of reduced BaGd1-xYbxMn2O5 (x = 0, 0.2 and 0.4) samples are rather different (see Fig. 3a–f), although the same synthesis condition was applied. As can be seen in Fig. 3a, the BaGdMn2O5 compound with majority of grains being around 5 μm in size, shows the biggest grain size among studied BaGd1-xYbxMn2O5 materials. However, when compared with the reported microstructure of BaGdMn2O5 in the literature [44], the present grain size in Fig. 3a is unfortunately larger, which in consequence affects the oxygen storagerelated properties. While in the case of BaGd0.8Yb0.2Mn2O5 material, particles with smaller size can be detected. The smallest grain size among investigated compounds is show in the BaGd0.6Yb0.4Mn2O5 sample. This indicates the increased content of Yb decreases the grain size and can possibly enhances the oxygen storage performance. The additional performed particle size analysis of BaGd1-xYbxMn2O5 materials (Fig. 3d–f) reveals the presence of large particles with 100 μm in BaGdMn2O5, and the smaller size of particles can also be detected. The substitution of Gd with Yb in BaGd0.8Yb0.2Mn2O5 oxide causes the

Fig. 4. XPS measurements of reduced BaGd1-xYbxMn2O5 compounds. 107

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Fig. 5. Isothermal reduction of BaGd1-xYbxMn2O6 at: a) 500 °C and b) 600 °C in 5 vol% H2 in argon.

Mn3+. The XPS data of all the reduced powder samples show the presence of the expected oxidation states of Mn2+ and/or Mn3+, and lack of Mn4+, which is of importance, as surface states of Mn cannot be easily predicted. This indicate that the reduced materials do not seem to be surface-oxidized. The data also correspond well with the TG measurements and determined oxygen content by iodometric titration, as well as the reported oxygen non-stoichiometry δ = 0 in BaGdMn2O5+δ (in 5% vol. H2/Ar) [21].

significantly improved by high energy milling, but at the same time further modification of the synthesis method (to obtain fine powders, e.g. at lower temperatures) may be also applied. Interestingly, addition of Yb in BaGd1-xYbxMn2O5+δ materials (apart from chemical modification, especially for BaGd0.6Yb0.4Mn2O5+δ) results also in the decreased content of larger secondary particles (aggregated and agglomerated), as is visible in Fig. 3. This is beneficial for the oxygen storage performance. At the same time, the primary particles (crystallites) are rather not modified by Yb introduction. Since the mechanism of the oxygen intercalation (during storage) involves surface adsorption and reduction, as well as bulk diffusion, larger specific surface and shorter (average) diffusion paths are highly beneficial. For the non-isothermal reduction and oxidation of BaGd1xYbxMn2O5+δ oxides, the heating/cooling rate with 5 °C/min was applied, and the results are presented in Fig. 6a–b and Table 3. The substitution of Gd with Yb (x = 0.4) decreases the non-isothermal reduction temperature (467 °C) and oxidation temperature (330 °C) when compared with the BaGdMn2O5+δ parent material. While in the case of BaGd0.8Yb0.2Mn2O5+δ compound, no evident enhancement can be detected, which indicates the doping effect of a small amount of Yb (x = 0.2) is almost negligible. Furthermore, the reduction and oxidation of BaGd1-xYbxMn2O5+δ oxides are reversible. In order to in-situ investigate the oxidation of BaGd1-xYbxMn2O5+δ oxides, the high temperature XRD measurements have been performed, and the results are discussed in the following section.

3.3. Oxygen storage properties As the reduction is the limiting process (oxidation is much faster in dozens of seconds), the present work focuses on the reduction measurements of studied materials. The results of the thermogravimetric measurements for BaGd1-xYbxMn2O5+δ materials conducted in isothermal conditions at 500 °C and 600 °C during reduction (in 5 vol% H2/Ar) cycles are presented in Fig. 5a and b. Oxygen storage-related parameters, calculated on the basis of these studies are gathered in Table 2. The increasing substitution of Gd with smaller Yb gradually decreases the theoretical oxygen storage capacity (OSC), due to the higher molar mass of Yb. The Yb doping in BaGd0.6Yb0.4Mn2O5+δ largely shortens the reduction time at 500 °C and 600 °C, which may be related with the smaller grain size and morphology of the powder and also the chemical composition [8]. The enhanced oxygen mobility of BaGd0.6Yb0.4Mn2O5+δ compound is particularly evident at 500 °C, and the reduction time is significantly shortened from 13.4 min (for BaGdMn2O5+δ) to 5.9 min (BaGd0.6Yb0.4Mn2O5+δ). The reduction time of BaGd1-xYbxMn2O5+δ oxides is considerably decreased at 600 °C (see Table 2), and the shortest reduction time of 2.3 min is recorded for BaGd0.6Yb0.4Mn2O5+δ. This proves that the Yb doping in BaGd1xYbxMn2O5+δ materials brings favorable influence on the oxygen storage properties and it enhances the oxygen diffusion, which indicates the chemical composition modification of BaLnMn2O5+δ series materials can be a good strategy to improve the current oxygen storage performance reported in the literature. Although the measured reduction time here is longer than the results measured in the previous work [9,22,28,29], the morphology of the investigated samples may still be significantly improved by high energy milling [30,31] and/or modified synthesis method [32], to markedly enhance the oxygen storage properties. The reported milling method [30,31] enhances the properties especially in terms of the oxygen release kinetics. Very fine powders of BaYMn2O5+δ have been successfully obtained by low-temperature firing (down to 900 °C) of precursors obtained from a wet-chemical route [32]. Since chemical composition and structure of the considered BaGd1-xYbxMn2O5+δ is similar to the parent BaLnMn2O5+δ, it can be expected that properties of the investigated samples can be also

3.4. In-situ studies of oxygen intake at high temperatures The in-situ studies of oxygen intake of BaGd0.6Yb0.4Mn2O5 sample have been performed by high temperature XRD measurements. Unfortunately, due to the technical limitations of Anton Paar HTK 1200N oven-chamber, it was not feasible to apply reducing atmosphere. Nevertheless, it is well known and measured by TG studies that the reduction process (even in a relatively strong reducing atmosphere of 5 vol% H2 in Ar at 500 °C [9,28]) takes much longer time (minutes), when comparing to the oxidation process (seconds). The origin of this behavior may stem from the fact that the oxidation process is associated with a strong heat release (for BaYMn2O5, exothermic ΔH ≈ −220 kJ mol−1 [46]), which causes a quick local over-heating of the material, accelerating the oxygen diffusion in the bulk. Moreover, the exothermic and endothermic effects have been observed in the TG measurements during the oxidation and reduction processes, respectively, and consequently it speeds up the oxygen intake and slows down the oxygen release. As can been seen in Fig. 7a and b, the oxidation of reduced BaGd0.6Yb0.4Mn2O5 to BaGd0.6Yb0.4Mn2O6 is recorded around 230 °C in air, 108

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3.30% 3.28% 3.13% 13.4 min 12.6 min 5.9 min

4.0 min 2.7 min 2.3 min

which is also observed in the TG measurement in Fig. 6b. The in-situ high temperature XRD measurement at each temperature point lasts around 1 h and in the condition of continuous air flow. Therefore, the presence of BaGd0.6Yb0.4Mn2O5.5 with small grains has not been observed, due to the fact the O5.5 composition (BaGd0.6Yb0.4Mn2O5.5) exists at a lower oxygen partial pressure [21]. The oxidation of BaGd0.6Yb0.4Mn2O5 to BaGd0.6Yb0.4Mn2O6 significantly decreases the unit cell parameters a and b, and in consequence decreases the cell volume V. While the influence on unit cell parameter c is trivial. This may indicate that the oxygen diffuses mainly along the (a, b) layered structure in BaGd0.6Yb0.4Mn2O5+δ during the oxygen intake process. The oxygen diffusion along c is less significant. The oxygen transport through the investigated material via the vacancy hopping mechanism likely involves the nearest-neighbor oxygen atom sites in the Gd0.6Yb0.4 layer (the vacancy-rich Gd0.6Yb0.4–O3 layer) [47]. The significant unit cell volume decrease occurring between 225 °C and 275 °C (see Fig. 7a) can be related with the oxidation process of BaGd0.6Yb0.4Mn2O5 to BaGd0.6Yb0.4Mn2O6. The oxidation process of BaGd0.6Yb0.4Mn2O5 is related with P4/nmm → P-1 space group change. Above 275 °C, the material is totally oxidized to BaGd0.6Yb0.4Mn2O6, which can be indicated by the linear unit cell volume change in function of temperature above this temperature. The phase transition of P-1 to P4/mmm in oxidized BaGd0.6Yb0.4Mn2O6 sample has also been observed around 450 °C (Fig. 7a). The unit cell parameters and volume data (in empty shapes with O6 note) of BaGd0.6Yb0.4Mn2O6 with P-1 structure at room temperature (25 °C) have also been added in Fig. 7a, and it shows a linear behavior (in function of temperature) with the high temperate P-1 data of BaGd0.6Yb0.4Mn2O6 (275–400 °C). This proves that oxidized BaGd0.6Yb0.4Mn2O6 material from room temperature to ~450 °C shows P-1 structure, while above 450 °C it presents P4/mmm space group. The P4/ mmm space group instead of P4/nmm has been applied to refine the high temperature XRD data above 450 °C, because the very high electrical conductivity of the sample at high temperatures indicates a significant delocalization of electrons near Fermi level, suggesting that Mn3+ and Mn4+ charge ordering does not take place. High temperature XRD data were also used for determination of the linear TEC of the considered materials. Generally, the measured values of BaGd0.6Yb0.4Mn2O5+δ material are moderate (data included in Fig. 7a), not exceeding 14.8(1) × 10−6 K−1. Interestingly, the calculated TEC values of BaGd0.6Yb0.4Mn2O5+δ material by high temperature XRD measurements are close to the commonly used solid electrolytes for SOFCs exhibit similar TEC values: La0.9Sr0.1Ga0.8Mg0.2O3–12.17 × 10−6 K−1, Zr0.85Y0.15O1.93 (8YSZ) – 10.8 × 10−6 K−1 and Ce0.8Gd0.2O1.9–12.5 × 10−6 K−1 [24]. This indicates that the thermal expansion of BaGd0.6Yb0.4Mn2O6 will not limit the wide potential application of this material, such as a good candidate of electrode materials for SOFCs [1,39,40]. 3.5. Transport properties of BaGd0.6Yb0.4Mn2O5+δ material

3.30% 3.28% 3.26% BaGdMn2O5+δ BaGd0.8Yb0.2Mn2O5+δ BaGd0.6Yb0.4Mn2O5+δ

3.26% 3.26% 3.11%

Change of weight on reduction/oxidation (5th cycle at 600 °C) Time of reduction (95% of total mass change, 5th reduction at 500 °C) Change of weight on reduction/oxidation (5th cycle at 500 °C) Theoretical change of weight Chemical composition

Table 2 Oxygen storage properties of BaGd1-xYbxMn2O5+δ oxides at 500 °C and 600 °C between air and 5% H2/Ar.

Time of reduction (95% of total mass change, 5th reduction at 600 °C)

K. Zheng

Fig. 8a shows the electrical conductivity of BaGd0.6Yb0.4Mn2O5+δ compound in air and in 5 vol% H2/Ar. Oxidized BaGd0.6Yb0.4Mn2O6 possesses much higher electrical conductivity, when comparing with the reduced one. At room temperature, the value of σ exceeds 0.1 Scm−1, and it gradually increases with the rise of temperature to around 200 °C. The activation energy Ea = 0.17 eV in the range of 30–200 °C is lower than that of BaErMn2O6 reported in the previous work [9]. Interestingly, between 200 and 550 °C, the activation energy decreases to Ea = 0.10 eV, while the electrical conductivity still increases. The same behavior was also observed in the BaErMn2O6 sample, and this may be related with the phase transition from charge/ orbital-ordered insulator COI(CE) to paramagnetic metal PM phase [9]. Further investigations are needed in order to determine the transition, as expected for BaLnMn2O6 with small Ln3+ cations [25,38]. The highest electrical conductivity for BaGd0.6Yb0.4Mn2O6 compound was recorded around 600 °C in air and the value reaches 100 Scm−1, which 109

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Fig. 6. Non-isothermal a) reduction in 5 vol% H2 in argon and b) oxidation in air of BaGd1-xYbxMn2O5+δ. Table 3 Oxygen storage properties of BaGd1-xYbxMn2O5+δ oxides on non-isothermal oxidation/reduction. Chemical composition

Theoretical change of weight

Average change of weight on nonisothermal oxidation/reduction

Temperature of reduction (95% of total mass change)

Temperature of oxidation (95% of total mass change)

BaGdMn2O5+δ BaGd0.8Yb0.2Mn2O5+δ BaGd0.6Yb0.4Mn2O5+δ

3.30% 3.28% 3.26%

3.26% 3.15% 3.10%

516 °C 515 °C 467 °C

346 °C 343 °C 330 °C

Fig. 7. a) Temperature dependence of normalized unit cell parameters and volume, together with calculated thermal expansion coefficients of reduced BaGd0.6Yb0.4Mn2O5 during heating in air, b) structural evolution of reduced BaGd0.6Yb0.4Mn2O5 during heating in air, data shown for selected angular range.

meets the requirement of cathode materials for SOFCs (~ 100 Scm−1). This behavior is also in a good agreement with the reported paramagnetic metal-type behavior for the Mn-contained compound [38]. It can also be stated that electron transfer between Mn3+ and Mn4+ states in BaGd0.6Yb0.4Mn2O6 compound is much easier. The recorded conductivity values for BaGd0.6Yb0.4Mn2O6 in this work, are much higher than those of BaLnMn2O5+δ materials reported in the literature [1,9,39,40]. At higher temperatures (above 600 °C), the measured conductivity of BaGd0.6Yb0.4Mn2O6 decreases, which is also a common feature of perovskite materials [41]. Additional equilibrated conductivity data have been collected during the oxidation process with enough long-time annealing at each temperature. The visible increase of conductivity at lower temperatures may be related with oxidation of material surface, which in consequence improves the conductivity. However, one has to notice that reaching full equilibrium at e.g. 100 °C is not possible, due to the kinetic reasons, but also due to intrinsic structural reasons (e.g. formation of the intermediate phase with

oxygen content equal 5.5). Since the material oxidizes during such the studies, the measure data can be only considered as relevant to the particular experimental conditions. An evident observed slop change of the conductivity behavior is around 300 °C, which may be related with the (bulk) oxidation process of the material. This has been observed in the in-situ XRD measurements (see Fig. 7a). The reduction of BaGd0.6Yb0.4Mn2O6 in 5 vol% H2 in argon significantly decrease the conductivity. During the reduction process, the conductivity starts a substantial decrease around 300 °C and the decrease stops at around 500 °C, which corresponds well with the reduction measurements of BaGd0.6Yb0.4Mn2O6 to BaGd0.6Yb0.4Mn2O5 studied by TG in Fig. 6a. The reduced BaGd0.6Yb0.4Mn2O5 sample presents much lower conductivity (almost 4 orders difference at room temperature) than that of oxidized one, and it shows the highest electrical conductivity 8.3 Scm−1 at 850 °C, which also meets the criteria of electrode materials for SOFCs (> 1 Scm−1). This may indicate that BaGd0.6Yb0.4Mn2O5+δ compound can also be a good candidate of 110

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Fig. 8. a) Electrical conductivity of BaGd0.6Yb0.4Mn2O5+δ as a function of temperature, b) oxygen transport coefficients D and k determination at 500 °C.

The increased content of Yb in BaGd1-xYbxMn2O5+δ compounds decreases the relative unit cell volume change between oxidized materials and reduced counterparts. Crystal structure with P4/nmm space group was observed in all BaGd1-xYbxMn2O5+δ (x = 0, 0.2 and 0.4) materials except for BaGd1-xYbxMn2O6 oxide (P-1 structure). The oxygen nonstoichiometry in the oxidized BaGd1-xYbxMn2O5+δ materials were determined as: 5.92(1), 5.97(1) and 5.93(1), respectively (for x = 0, 0.2 and 0.4 samples), and the non-stoichiometry of fully reduced samples was also calculated based on the TG data, and the oxygen content was found to be close to 5.0 (slightly below). The increased content of Yb in BaGd1-xYbxMn2O5+δ improves the materials' morphology. The Mn2+ and/or Mn3+ oxidation states were present in the reduced BaGd1xYbxMn2O5 materials, confirmed by the XPS studies and oxygen nonstoichiometry determination. The Yb doping in BaGd1-xYbxMn2O5+δ significantly decreases the reduction time (at 500 and 600 °C during isothermal measurements) and oxidation/reduction temperature during non-isothermal runs. BaGd0.6Yb0.4Mn2O5+δ sample shows the best oxygen storage-related properties among the investigated materials. The oxygen in-situ intake during oxidation in air related with P4/nmm → P-1 space group change is between 225 and 275 °C, the nature of which may indicate the oxygen diffusion in this material mainly occurs along the vacancy-rich Gd0.6Yb0.4–O3 layer structure. For oxidized BaGd0.6Yb0.4Mn2O6, the P-1 → P4/mmm phase transition is observed at around 450 °C. The calculated TEC values (not exceeding 14.8(1) × 10−6 K−1) from high temperature XRD data are moderate. The recorded high electrical conductivities (100 Scm−1 in air at 600 °C, and 8.3 Scm−1 in 5 vol% H2/Ar at 850 °C) and fast oxygen diffusion coefficients (D = (5.6–6.3) × 10−6 cm2 s−1 and k = (4.0–5.0) × 10−4 cms−1 at 500 °C) suggest that BaGd0.6Yb0.4Mn2O5+δ can have wide potential applications such as: good candidate of electrode materials for SOFCs.

electrode materials for SOFCs [1,39,40]. The energy activation of BaGd0.6Yb0.4Mn2O5 in 200–600 °C is about 0.31 eV, and this is much higher than those recorded in the air. The conductivity recorded during the oxidation of BaGd0.6Yb0.4Mn2O5 to BaGd0.6Yb0.4Mn2O6, increases gradually with visible changes around 225 °C and 400 °C, which may be related with the phase transitions recorded in Fig. 7a. The determination of oxygen chemical diffusion D and surface exchange constant k of BaGd0.6Yb0.4Mn2O6 has also been conducted by mass relaxation technique in TG measurements at 500 °C [35,36]. The relaxtion measurments have been performed for samples with two different thicknesses (thin and thick samples). Oxygen partial pressure change was between 0.1 and 0.01 atm. It is known that the oxygen partial pressure change (p1/p2) is very essential in relaxation measurements, and it is reported that the oxygen partial pressure switches with p1/p2 > 20, the assumption of linear exchange kinetics is no longer valid [48,49]. Moreover, the transport coefficients D and/or k may be dependent or independent on the oxygen partial pressure [50–53]. Therefore, the oxygen partial pressure change should be as small as possible based on the specific testing system. However, one has also to note that a small step size, unfortunately, can also cause a poorer signal-to-noise ratio. In this work, for the relaxation measurements the oxygen partial pressure change was optimally chosen between 0.1 and 0.01 atm. The reactor (TG chamber) flush time is evaluated to be < 1 s, which maximally decrease the influence of the flush time on the data. One dimension geometry was used, as the samples were flat in shape, with thickness of 0.21 mm and 0.60 mm (other dimensions at leat 10 times larger). As can been seen from Fig. 8b, the fitting quality is very good, and fitted curves match very closely the data. The determined D and k values for the two samples with different thickness are very cloese to each other. Furthermore, the calculated dimensionless parameters lk L = D were L = 0.8 (thin sample) and L = 2.1 (thick sample), which both fullfill the requriment of simultanious determination of D and k: lk 0.03 ≤ L = D ≤ 30 [54]. This can confirm that the obtained D and k are reliable. As can be seen in Fig. 8b, high oxygen transport coefficients D = (5.6–6.3) × 10−6 cm2 s−1 and k = (4.0–5.0) × 10−4 cms−1 have been determined at 500 °C. The measured D and k values are comparable with those of cathode materials for SOFCs [35,41]. The measured oxygen transport data can be a good indicator of enhanced oxygen mobility in BaGd0.6Yb0.4Mn2O6 material.

Acknowledgments The author would like to thank Foundation for Polish Science (FNP) for the support, and Jacek Jagusztyn for his help during the synthesis of materials as well as prof. Konrad Świerczek for fruitful discussions. References [1] S. Sengodan, S. Choi, A. Jun, T.H. Shin, Y.-W. Ju, H.Y. Jeong, J. Shin, J.T.S. Irvine, G. Kim, Nat. Mater. 14 (2) (2015) 205–209. [2] F. Tonus, M. Bahout, V. Dorcet, G.H. Gauthier, S. Paofai, R.I. Smithd, S.J. Skinner, J. Mater. Chem. A 4 (2016) 11635–11647. [3] T. Motohashi, T. Ueda, Y. Masubuchi, M. Takiguchi, T. Setoyama, K. Oshima, S. Kikkawa, Chem. Mater. 22 (2010) 3192–3196. [4] Y. Lu, H. Zhao, X. Chang, X. Du, K. Li, Y. Ma, S. Yi, Z. Du, K. Zheng, K. Swierczek, J. Mater. Chem. A 4 (27) (2016) 10454–10466. [5] T. Motohashi, M. Kimura, Y. Masubuchi, S. Kikkawa, J. George, R. Dronskowski,

4. Conclusions In this work, the favorable influence of Yb doping in BaGd1double perovskites has been recorded on the crystal structure, oxygen nonstoichiometry, microstructure and oxidation states of Mn, oxygen storage properties, oxygen in-situ intake at high temperatures, and electrical conductivity, as well as oxygen diffusion. xYbxMn2O5+δ

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[30] A. Klimkowicz, K. Świerczek, T. Yamazaki, A. Takasaki, Solid State Ionics 298 (2016) 66–72. [31] K. Świerczek, A. Klimkowicz, A. Takasaki, K. Zheng, T. Yamazaki, “Method of manufacturing of highly-efficient oxygen storage materials”, The Application Registered in the European Patent Office, 2015 Under no. EP15186419.6. [32] T. Motohashi, T. Ueda, Y. Masubuchi, S. Kikkawa, J. Ceram. Soc. Jpn. 119 (2011) 894–897. [33] A. Olszewska, Z. Du, K. Świerczek, H. Zhao, B. Dabrowski, J. Mater. Chem. A 6 (2018) 13271–13285. [34] B.H. Toby, J. Appl. Crystallogr. 34 (2001) 210–213. [35] K. Zheng, A. Gorzkowska-Sobaś, K. Świerczek, Mater. Res. Bull. 47 (2012) 4089–4095. [36] K. Zheng, K. Świerczek, Solid State Ionics 323 (2018) 157–165. [37] J. Crank, The Mathematics of Diffusion, 2nd ed., Oxford University Press, New York, 1975. [38] Y. Ueda, T. Nakajima, J. Phys. Condens. Matter 16 (2004) S573–S583. [39] O.L. Pineda, Z.L. Moreno, P. Roussel, K. Świerczek, G.H. Gauthier, Solid State Ionics 288 (2016) 61–67. [40] S. Choi, S. Sengodan, S. Park, Y.W. Ju, J. Kim, J. Hyodo, H.Y. Jeong, T. Ishihara, J. Shin, G. Kim, J. Mater. Chem. A 4 (2016) 1747–1753. [41] C. Kuroda, K. Zheng, K. Świerczek, Int. J. Hydrog. Energy 38 (2013) 1027–1038. [42] K. Jeamjumnunja, W. Gong, T. Makarenko, A.J. Jacobson, J. Solid State Chem. 230 (2015) 397–403. [43] M. Karppinen, H. Okamoto, H. Fjellvag, T. Motohashi, H. Yamauchi, J. Solid State Chem. 177 (2004) 2122–2128. [44] E. Tsuji, T. Motohashi, H. Noda, Y. Aokic, H. Habazaki, J. Phys. Chem. C 122 (13) (2018) 7081–7087. [45] National Institute of Standards and Technology XPS database, http://srdata.nist. gov/xps/Default.aspx. [46] M. Gilleßen, M. Lumeij, J. George, R. Stoffel, T. Motohashi, S. Kikkawa, R. Dronskowski, Chem. Mater. 24 (2012) 1910–1916. [47] R.A. Cox-Galhotra, A. Huq, J.P. Hodges, J.-H. Kim, C. Yu, X. Wang, A.J. Jacobson, S. McIntosh, J. Mater. Chem. A 1 (2013) 3091–3100. [48] S. Wang, A. Verma, Y.L. Yang, A.J. Jacobson, B. Abeles, Solid State Ionics 140 (2001) 125–133. [49] B. Hu, C. Xia, Asia Pac. J. Chem. Eng. 11 (2016) 327–337. [50] W. Preis, E. Bucher, W. Sitte, J. Power Sources 106 (2002) 116–121. [51] P.-M. Geffroy, L. Guironnet, H.J.M. Bouwmeester, T. Chartier, J.-C. Grenier, J.M. Bassatd, J. Eur. Ceram. Soc. 39 (2019) 59–65. [52] I. Yasuda, T. Hikita, J. Electrochem. Soc. 141 (1994) 1268–1273. [53] J.E. ten Elshof, M.H.R. Lankhorst, H.J.M. Bouwmeester, Solid State Ionics 99 (1997) 15–22. [54] M.W. den Otter, H.J.M. Bouwmeester, B.A. Boukamp, H. Verweij, J. Electrochem. Soc. 148 (2) (2001) J1–J6.

Chem. Mater. 28 (12) (2016) 4409–4414. [6] J. Vieten, B. Bulfin, F. Call, M. Lange, M. Schmucker, A. Francke, M. Roeb, C. Sattler, J. Mater. Chem. A 4 (2016) 13652–13659. [7] Q. Song, W. Liu, C.D. Bohn, R.N. Harper, E. Sivaniah, S.A. Scottc, J.S. Dennis, Energy Environ. Sci. 6 (2013) 288–298. [8] A. Klimkowicz, K. Świerczek, A. Takasaki, J. Molenda, B. Dabrowski, Mater. Res. Bull. 65 (2015) 116–122. [9] K. Świerczek, A. Klimkowicz, K. Zheng, D. Dabrowski, J. Solid State Chem. 203 (2013) 68–73. [10] Y. Nagai, T. Yamamoto, T. Tanaka, S. Yoshida, T. Nonaka, T. Okamoto, A. Suda, M. Sugiura, Catal. Today 74 (2002) 225–234. [11] M. Boaro, F. Giordano, S. Recchia, V.D. Santo, M. Giona, A. Trovarelli, Appl. Catal. B Environ. 52 (2004) 225–237. [12] M. Karppinen, H. Yamauchi, S. Otani, T. Fujita, T. Motohashi, Y.H. Huang, M. Valkeapää, H. Fjellvåg, Chem. Mater. 18 (2006) 490–494. [13] O. Parkkima, H. Yamauchi, M. Karppinen, Chem. Mater. 25 (2013) 599–604. [14] M. Hervieu, A. Guesdon, J. Bourgeois, E. Elkaїm, M. Poienar, F. Damay, J. Rouquette, A. Maignan, C. Martin, Nat. Mater. 13 (2013) 74–80. [15] S. Remsen, B. Dabrowski, Chem. Mater. 23 (2011) 3818–3827. [16] C. Abughayada, B. Dabrowski, S. Kolesnik, D.E. Brown, O. Chmaissem, Chem. Mater. 27 (18) (2015) 6259–6267. [17] T. Motohashi, Y. Hirano, Y. Masubuchi, K. Oshima, T. Setoyama, S. Kikkawa, Chem. Mater. 25 (2013) 372–377. [18] M. Machida, K. Kawamura, T. Kawano, D. Zhang, K. Ikeue, J. Mater. Chem. 16 (2006) 3084–3090. [19] K. Ikeue, M. Eto, D. Zhang, T. Kawano, M. Machida, J. Catal. 248 (2007) 46–52. [20] A. Klimkowicz, K. Świerczek, K. Zheng, D. Wallacher, A. Takasaki, J. Mater. Sci. 52 (2017) 6476–6485. [21] K. Jeamjumnunja, W. Gong, T. Makarenko, A.J. Jacobson, J. Solid State Chem. 239 (2016) 36–45. [22] A. Klimkowicz, K. Zheng, G. Fiołka, K. Świerczek, Chemik 67 (12) (2013) 1199–1206. [23] G. King, P.M. Woodward, J. Mater. Chem. 20 (2010) 5785–5796. [24] K. Zheng, K. Świerczek, J. Eur. Ceram. Soc. 34 (2014) 4273–4284. [25] T. Nakajima, H. Kageyama, H. Yoshizawa, Y. Ueda, J. Phys. Soc. Jpn. 71 (12) (2002) 2843–2846. [26] K. Świerczek, A. Klimkowicz, A. Niemczyk, A. Olszewska, T. Rząsa, J. Molenda, A. Takasaki, Funct. Mater. Lett. 7 (6) (2014) 1440004. [27] A. Klimkowicz, K. Świerczek, T. Rząsa, A. Takasaki, B. Dabrowski, Solid State Ionics 288 (2016) 43–47. [28] A. Klimkowicz, K. Świerczek, K. Zheng, M. Baranowska, A. Takasaki, B. Dabrowski, Solid State Ionics 262 (2014) 659–663. [29] A. Klimkowicz, K. Zheng, G. Fiołka, K. Świerczek, Materiały Ceramiczne/Ceramic Materials 65 (1) (2013) 92–96.

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