A novel layered perovskite as symmetric electrode for direct hydrocarbon solid oxide fuel cells

Journal of Power Sources 342 (2017) 313e319

Contents lists available at ScienceDirect

Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour

A novel layered perovskite as symmetric electrode for direct hydrocarbon solid oxide fuel cells Ling Zhao a, Kongfa Chen b, Yuanxu Liu c, Beibei He a, * a

Department of Material Science and Chemistry, China University of Geosciences, Wuhan, 430074, China College of Materials Science and Engineering, Fuzhou University, Fuzhou, Fujian 350108, China c Department of Materials Science and Engineering, University of Science and Technology of China, Hefei, 230026, China b

h i g h l i g h t s

g r a p h i c a l a b s t r a c t


A novel layered PrBaMn1.5Fe0.5O5þd (PBMFO) perovskite is synthesized.
The layered PBMFO exhibits high electrical conductivity.
Cells using symmetrical PBMFO electrodes generate competitive performance.
PBMFO anode shows good redox stability and high carbon tolerance.

a r t i c l e i n f o

a b s t r a c t

Article history: Received 27 September 2016 Received in revised form 12 December 2016 Accepted 16 December 2016

Layered perovskite oxides are well known to possess significant electronic, magnetic and electrochemical properties. Herein, we highlight a novel layered perovskite PrBaMn1.5Fe0.5O5þd (PBMFO) as electrodes of symmetrical solid oxide fuel cells (SSOFCs). The layered PBMFO shows high electrical conductivity of 112.5 and 7.4 S cm1 at 800  C in air and 5% H2/Ar, respectively. The single cell with PBMFO symmetric electrodes achieves peak power density of 0.54 W cm2 at 800  C using humidified hydrogen as fuel. Moreover, PBMFO electrodes demonstrate good redox stability and high coking tolerance against hydrocarbon fuel. © 2016 Elsevier B.V. All rights reserved.

Keywords: Solid oxide fuel cells Hydrocarbon Symmetric electrodes Layered perovskite

1. Introduction Solid oxide fuel cells (SOFCs), a type of electrochemical devices converting chemical energy directly to electric power, are regarded as the feasible and environmentally friendly technology for energy supply because of their fuel flexibility and high energy conversion efficiency properties [1e6]. The single cell of SOFCs is usually a sandwich structure, consisting of a dense electrolyte layer, a porous

* Corresponding author. E-mail address: babyfl[email protected] (B. He). http://dx.doi.org/10.1016/j.jpowsour.2016.12.066 0378-7753/© 2016 Elsevier B.V. All rights reserved.

anode and cathode. The porous anode and cathode is exposed to a reducing and an oxidizing atmosphere, respectively. Recently, symmetrical SOFCs (SSOFCs), where anode and cathode materials are the same, have received increasing interest owing to the fact that such symmetrical configuration can substantially simplify the manufacture and save the cost of SOFCs [7e10]. Besides, carbon deposition and/or sulfur poisoning can be overcome or suppressed because of the possibility of reversing the supplied gases, oxidizing the possible carbon deposit and/or sulfur species, and consequently recovering the loss of output power density [7]. In order to perform efficiently for a long lifetime, those electrode materials of SSOFCs should meet the following requirements:

314

L. Zhao et al. / Journal of Power Sources 342 (2017) 313e319

robust structural stability against redox, sufficient electrical conductivity in reducing and oxidizing atmospheres, and excellent electrocatalytic activity for fuel oxidation and oxygen reduction. However, only a few materials have been proposed as symmetrical electrodes because of their strict restrictions. Mixed ionicelectronic conductor (MIEC) perovskite oxides are considered as the most promising type of symmetrical electrodes of SSOFCs, such as (La1xSrx)0.9Cr0.5Mn0.5O3d, Sr2Fe1.5Mo0.5O6d, La0.4Sr0.6Ti1La0.8Sr0.2Sc0.2Mn0.8O3d, SrFe0.75Zr0.25O3d, yCoyO3±d, La0.4Sr0.6Co0.2Fe0.7Nb0.1O3d, LaSr2Fe2CrO9d and LaCo0.3Fe0.67Pd0.03O3d et al. [8,11e17]. Another advantage of these symmetrical perovskite electrodes is their high tolerance against carbon deposition and sulfur poisoning. The performances of SSOFCs using perovskite electrodes, however, are still inadequate for practical application primarily due to the relatively low catalytic activity. Therefore, it is of great scientific and practical significance to develop SSOFCs considering highly active and durable symmetrical electrode materials. Recently, A-site ordered LnBaMn2O5þd (LnBMO, Ln ¼ Pr and Nb) perovskites have been reported as highly efficient electrodes of SOFCs [18e22]. The layered LnBMO perovskites can be described with the following stacking sequence [LnOx]e[MnO2]e[BaO]e [MnO2]e[LnOx] along the c axis. Such layered perovskite structure can decrease the bond strength in [AO] layer and create the disorder-free channel for oxygen species transfer, and thereby promising the diffusion of oxygen vacancy. Further, it was demonstrated that PrBaMn2O5þd (PBMO) anode had good redox stability with high tolerance against carbon deposition and sulfur poisoning [18]. It is generally accepted that the crystal structures of perovskites are quite stable against amounts of doping on A and/or B sites due to their high tolerance against crystal structure distortion [23e26]. Therefore, several progresses have been made in developing perovskites by partial substitution, in order to obtain the desired electrical, catalytic, and stable properties. For example, the Mo-doped PBMO anode exhibited the promoted electrocatalytic activity for fuel oxidation of H2 and CH4, which might be attributed to the creation of oxygen vacancies by Mo doping [27]. In the case of Ca-doped PBMO, the layered PrBa0.8Ca0.2Mn2O5þd (PBCMO) symmetrical electrodes demonstrated outstanding electrochemical activity and sufficient stability in various hydrocarbon fuels [28]. As known, the electrocatalytic activities of perovskites are primarily determined by B-site cations and oxygen vacancies [29e31]. It was reported that the transition metals(Fe, Mn and Co)-rich perovskites provided high electrocatalytic activity in relation to direct hydrocarbon oxidation [30,32e35]. In this study, we reported a novel layered perovskite with Fe doping as symmetric electrodes of SSOFCs. The aim of this study was to demonstrate the feasibility of layered perovskite PrBaMn1.5Fe0.5O5þd (PBMFO) as symmetrical electrodes in hydrogen and hydrocarbon fuels. Layered perovskite PBMFO demonstrated high electrical conductivity, enhanced electrocatalytic activity and good durability in reducing and oxidizing atmospheres, indicating its appreciated potential as competitive symmetrical electrodes of SSOFCs.

agents of citric acid and EDTA were then introduced at a molar ratio of 1: 1.5: 1 (metal ions: citric acid: EDTA). The pH value of such mixture was adjusted to approximately 7 with the assistance of NH3$H2O. After stirring for one hour, the resulting solution was treated on heating stage until self-ignition. The as-prepared powder was then sintered in air at 600  C for 3 h and subsequently at C 1100 for 4 h to obtain A-site dis-ordered Pr0.5Ba0.5Mn0.75Fe0.25O3d perovskite. The A-site ordered PrBaMn1.5Fe0.5O5þd (PBMFO) phase was formed by annealing Pr0.5Ba0.5Mn0.75Fe0.25O3d in 5% H2/Ar at 800  C for 5 h La0.8Sr0.2Ga0.8Mg0.2O3d (LSGM) powder was synthesized via solid state reaction (SSR) way [39]. Stoichiometric amounts of La2O3, SrCO3, Ga2O3, and MgCO3 were mixed and ball milled with ethanol for 48 h. The mixture after drying was sintered at 1300  C for 5 h. 2.2. Cell preparation The as-prepared LSGM powder was dry-pressed at 300 MPa,

2. Experimental 2.1. Powder synthesis PrBaMn2O5þd (PBMO) and PrBaMn1.5Fe0.5O5þd (PBMFO) powders were synthesized by a citric acid-EDTA method [36e38]. Take the preparation of Pr0.5Ba0.5Mn0.75Fe0.25O3d for example, Pr(NO3)3$6H2O, Ba(NO3)2, Mn(NO3)2$4H2O and Fe(NO3)3$9H2O metal nitrates (Sinopharm Chemical Reagent Co., Ltd, analytical grade) were dissolved in deionized water. The parallel complexing

Fig. 1. X-ray diffraction patterns (XRD) patterns of (a) Pr0.5Ba0.5Mn0.75Fe0.25O3d powders after annealed at 1100  C in air for 4 h; (b) PBMFO powders after annealed at 800  C in 5% H2/Ar for 5 h; (c) PBMFO powders re-oxidized in air at 800  C for 4 h and (d) Refined diffraction patterns of PBMFO powders after annealed at 800  C in 5%H2-Ar for 5 h.

L. Zhao et al. / Journal of Power Sources 342 (2017) 313e319

315

subsequently sintered in air at 1450  C for 5 h to obtain LSGM disk. The thickness of dense LSGM disk was controlled at about 0.5 mm by polishing. The cell configuration of PBMFOjLSGMjPBMFO was used for symmetric and single cells testing. To prepare the PBMFO electrode, a slurry consisting of PBMFO and terpineol as a weight ratio of 1:2 was screen-printed [40e42] onto both side of the LSGM surface, and then sintered at 1100  C for 4 h to form symmetrical electrodes. The obtained symmetrical electrodes were about 50e60 mm in thickness. The effective area of symmetrical electrodes was 0.24 cm2.

the Van der Pauw method [44,45] from 600 to 800  C. Electrochemical workstation (Reference 3000, Gamry) was used to characterize the current-voltage curves and impedance spectra. For the single cell test, the cells were sealed on the top of alumina tube by glass sealants. Pt paste and Ag wires were used for current collecting. Humidified (~3% H2O) hydrogen and methane with flow rate of 50 mL min1 were applied as the fuels, and ambient air played as the oxidant.

2.3. Characterization

3.1. Phase structure

Phase structure was identified using X-ray diffraction (XRD, Bruker AXS D8-Focus) technique with CuKa radiation. The XRD Rietveld structural refinement was conducted via Fullporf software [43]. The microstructure was carried out by field emission scanning electron microscopy (FESEM, SU-8010) and high resolution transmission electron microscopy (HRTEM, Tecnai G2 F20 U-TWIN). Thermogravimetric analysis (TGA, SDT Q600) were investigated from room temperature to 800  C with a heating rate of 10  C min1 under H2 atmosphere. Hydrogen temperature programmed reduction (H2-TPR) was conducted by a micromerit autochem (II 2920) co-operated with a thermal conductivity detector. Each pretested sample was annealing in He at 1000  C for 0.5 h. The 5% H2/ Ar with the flow rate of 20 mL min1 was applied during TPR test. The heating rate of environmental temperature was 5  C min1. Thermal expansion coefficients (TECs) were studied on a dilatometer (SHIMADZU50) from room temperature to 900  C at a heating rate of 10  C min1. Dense bar samples as identified by the Archimedes method (~14.00  5.05  0.60 mm) were prepared for the study of electrical conductivities using a dc four-probe technique in oxidizing and reducing (5% H2/Ar) atmospheres, respectively. The electrical conductivity of the porous PBMO and PBMFO anode were measured by

A-site ordered PrBaMn1.5Fe0.5O5þd (PBMFO) could be obtained via a two-step process. First, dis-ordered Pr0.5Ba0.5Mn0.75Fe0.25O3d was synthesized in air at 1100  C for 4 h via a citric acid-EDTA approach. Then, the as-synthesized powder was annealing in 5% H2/Ar atmosphere at 800  C for 5 h to receive a layered perovskite structure. As shown in Fig. 1a, the fresh Pr0.5Ba0.5Mn0.75Fe0.25O3d synthesized in air at 1000  C exhibited a mixture of cubic (C) and hexagonal (H) structures. In comparison, only a tetragonal (T) structure was obtained after annealing in reducing atmosphere (Fig. 1b), implying that the possible phase change occurred during the reduction process. The layered perovskite structure of PBMFO was further identified by the XRD Rietveld refinement technique (Fig. 1d). The observed diffraction peaks of PBMFO corresponded to a layered tetragonal structure with space group of P4/mmm. The cell parameter c of 7.7612(17) Å was nearly twice as large as a of 3.8984(7) Å because of the ordering of Pr3þ and Ba2þ. Schematic representation of A-site dis-ordered cubic phase and A-site ordered tetragonal phase was described in Fig. S1. The similar phenomena of phase change were confirmed in PrBaMn2O5þd (PBMO) and PrBa0.8Ca0.2Mn2O5þd (PBCMO) as well [18,28]. Candidate materials as SSOFCs electrodes should keep robust structure under reducing and oxidizing atmospheres. In this regard, the layered PBMFO was

3. Results and discussion

Fig. 2. TEM analysis. (a) Bright-field TEM, (b) High-resolution TEM image and corresponding fast-Fourier transformed pattern with zone axis (Z. A.) ¼ [100], (c) High-angle annular dark-field scanning TEM image of Pr0.5Ba0.5Mn0.75Fe0.25O3d viewed in the [100] direction with d-spacing 001; (d) Bright-field TEM image, (e) High-resolution TEM image and corresponding fast-Fourier transformed pattern with (Z. A.) ¼ [100], (f) High-angle annular dark-field scanning TEM image of A-site layered PrBaMn1.5Fe0.5O5þd aligned along the [100] direction with d-spacing 001.

316

L. Zhao et al. / Journal of Power Sources 342 (2017) 313e319

100 Weight Percentage (%)

measured by re-annealing in air after the annealing in 5% H2/Ar (Fig. 1c). The layered structure was maintained without any secondary phase, suggesting that layered PBMFO was stable under both oxidizing and reducing atmospheres. Fig. 2 shows HRTEM views of dis-ordered Pr0.5Ba0.5Mn0.75Fe0.25O3d and ordered PBMFO samples. It was seen that the edges of air-prepared Pr0.5Ba0.5Mn0.75Fe0.25O3d particles were smooth (Fig. 2a and b). However, the morphology of PBMFO particles changed obviously after reduction. The layered PBMFO exhibited clear edges and facets (Fig. 2d and e). Moreover,

(a)

98

96 PBMO PBMFO

94

92

100

200

300

400

500

600

700

800

o

Temperature ( C)

(b)

Intensity (a.u.)

H2-TPR

PBMFO

PBMO

100

200

300

400

500

600

700

800

o

Temperature ( C) Fig. 3. (a) TGA analysis of PBMO and PBMFO powders in H2; (b) H2-TPR patterns of PBMO and PBMFO powders.

Fig. 4. Temperature dependence of electrical conductivities of PBMO and PBMFO, measured in air and 5% H2/Ar, compared to other two materials (LSCM [2] and SMMO [1]) measured in the same condition.

Fig. 5. Interfacial polarization resistance of the PBFMO electrode under air and 5% H2/ Ar, as a function of temperature, (b) Impedance spectra of the half-cells with pure PBMO and PBMFO symmetrical electrode measured under open circuit conditions in air at 800  C, (c) Impedance spectra of the half-cells with pure PBMO and PBMFO symmetrical electrodes measured in 5% H2/Ar at 800  C.

L. Zhao et al. / Journal of Power Sources 342 (2017) 313e319

the A-site ordered perovskite structure was confirmed by showing an additional spot in the Fourier transformed pattern (Fig. 2e), compared with the A-site disordered sample (Fig. 2b). It was also confirmed by the HRTEM image in Fig. 2f, implying a change of periodic contrast due to stacking sequence [PrOx]e[(Mn,Fe)O2]e [BaO]e[(Mn,Fe)O2]e[PrOx] in the lattice structure, compared with dis-ordered one in Fig. 2c. The lattice diffraction fringes was about 0.77 nm in Fig. 2f, which agreed well with the (001)PBMFO crystal face calculated from Rietveld refinement. The phase change associated with the reduction process of PBMO and PBMFO is assessed by thermogravimetric measurement in a reducing atmosphere (Fig. 3a). The small peak occurring around 300e400  C for both samples indicated the possible phase change process as well. In addition, the total weight loss of PBMFO was approximately 6.10%, which was larger than that of PBMO about 4.76%. The difference of weight loss proved that the substitution of Mn to Fe could create more oxygen vacancies. Fig. 3b presents the temperature programmed reduction (TPR) profile of PBMO and PBMFO samples. The hydrogen consumption peak of PBMFO shifted to a lower temperature range (~350  C), compared to PBMO sample (~400  C). The TPR results indicated that the PBMFO oxide contained more labile oxygen than the unprompted PBMO. This liability of the oxygen was expected to promote the redox capacity of PBMFO and thus affected the formation of oxygen vacancies which contributed to its electrochemical performance. 3.2. Electrochemical property Fig. 4 presents the electrical conductivities of PBMFO and PBMO as a function of temperature from 300  C to 800  C in air and 5% H2/ Ar, respectively. The PBMFO exhibited the relatively high electrical conductivities under air and 5% H2/Ar atmospheres. The electrical conductivities of PBMFO were higher than those of PBMO, which achieved a maximum of 112.5 and 7.4 S cm1 at 800  C in air and 5% H2/Ar, meeting the basic requirement of electrode materials for SOFCs [18]. Besides, it was clearly seen that PBMFO developed in this study provided much higher electrical conductivity than other

317

competitive ceramic oxide anodes (La0.75Sr0.25Cr0.5Mn0.5O3 [2] and Sr2MgMoO6d [1]), suggesting that PBMFO could be a highperforming anode material in terms of electrical conductivity. In layered PBMFO, XPS results (Fig. S3 and Table S1) indicated the presence of Fe2þ/Fe3þ and Mn3þ/Mn4þ pairs. The mobile electronic holes, Mn
Mn and Fe
Fe , which probably made a significant contribution to electronic conduction. After reduction, electronic holes in PBMFO reduced, resulting in the loss of electrical conductivity (Fig. 4 and Fig. S4). The electro-catalytic activities of PBMO and PBMFO electrodes toward the oxygen reduction and hydrogen oxidation were performed based on symmetrical cells under different atmospheres. Obviously, the cross-sectional microstructures of PBMO and PBMFO electrodes displayed quite similar (Fig. S5). The particles of PBMO and PBMFO electrodes were fine with the similar average sizes of ~0.2 mm, contacting well with LSGM electrolyte. It is well known that chemical and thermal compatibility between electrode and electrolyte is critical for contact in interface. The average thermal expansion coefficient (TEC) of PBMO and PBMFO was 12.8  106 K1 and 13.1  106 K1, respectively (Fig. S6), which was close to that of LSGM electrolytes (12e13  106 K1) [46,47]. In addition, XRD analysis implied no obvious secondary phase formed in the mixture of PBMFO-LSGM after heat-treatment at 1100  C (Fig. S7). The results demonstrated that the PBMFO electrode had a good chemical and thermal compatibility with the LSGM electrolyte. The electrical conductivity of PBMO and PBMFO was 1.25 and 0.50 S cm1 at 800  C in 5% H2/Ar, respectively, which was lower than the dense PBMO and PBMFO sample due to the presence of porosity (Fig. S8). Importantly, the electrode activity was enhanced after Fe doping through reducing the area-specificresistance (ASR) by 33% (from 0.30 U cm2 to 0.22 U cm2) in air and by 23% (from 0.78 U cm2 to 0.68 U cm2) in 5% H2/Ar at 800  C (Fig. 5). Therefore, the lower ASR demonstrated that Fe doping in Bsite of PBMO played the positive effect on the electrochemical activity and PBMFO could promise high electrocatalytic activity as cathode and anode of SOFCs. In this study, the Fe-doping and the intrinsic oxygen vacancies affected by the Fe-doping acted the

Fig. 6. SEM images of (a) cross-sectional view of the single cell after measurement, (b) PBMFO anode and electrolyte interface. (c) and (d) Morphology of PBMFO anode layer in different magnification.

318

L. Zhao et al. / Journal of Power Sources 342 (2017) 313e319

crucial roles in the electrode performance. Further studies are required to reveal the fundamental of these influencing factors. 3.3. Cell performance The cross-sectional microstructure of the electrolyte supported single cell using symmetrical PBMFO electrodes after testing is shown in Fig. 6. The single cell consisted of a ~520 mm dense LSGM electrolyte layer, a ~60 mm porous PBMFO anode and cathode. Prior

Fig. 7. (a) Typical IeV and I-P curves of the symmetrical PBMFO electrodes using hydrogen (3 vol.% H2O) fuels and ambient air as the oxidant. (b) single cell performance directly using wet CH4 as fuel (3 vol.% H2O). (c) Short term stability under a constant voltage load of 0.5 A cm2 at 750  C in humidified H2 and CH4, respectively.

to conducting the test of single cells, both electrodes were annealed in 5% H2/Ar to obtain layered structure. Fig. 7 displays the electrochemical performance of cells with construction of PBMFOjLSGMjPBMFO using H2 and CH4 as fuels. In the case of humidified H2 (3% H2O) as fuel, the maximum power densities of the cell using PBMFO symmetrical electrodes were 0.54 and 0.23 W cm2 at 800 and 700  C, respectively (Fig. 7a). These power densities were higher than prior reported SSOFCs, such as PrBaMn2O6d electrode on a ~310 mm thick LSGM electrolyte (0.57 W cm2 at 850  C) [18], Sr2Fe1.5Mo0.5O6d electrode on a 265 mm thick LSGM electrolyte (0.51 W cm2 at 800  C) [8], SrFe0.75Zr0.25O3/50 wt% Ce0.8Gd0.2O1.9 electrode on a ~400 mm thick LSGM electrolyte (0.43 W cm2 at 800  C) [15], and La0.8Sr0.2Sc0.2Mn0.8O3þd electrode on a 300 mm thick scandiumstabilized zirconia electrolyte (0.31 W cm2 at 900  C) [23]. The advanced electrochemical performance of PBMFO electrode could be attributed to its sufficient electrical conductivity and high electrocatalytic activity for fuel oxidation and oxygen reduction. In the case of using humidified CH4 (3% H2O) as fuel, the peak power densities of the cell using PBMFO electrodes were 0.34 and 0.17 W cm2 at 850 and 800  C, respectively, as presented in Fig. 7b. Besides, such PBMFO electrodes would greatly expand its application in anode supported fuel cells, when several technical strategies, such as introducing high electrically metallic catalysts into ceramic anode, are required to improve the total anode conductivity and catalytic activity [18,28,48]. On the other side, a thin LSGM film electrolyte would significantly reduce ohmic resistance of the cell and thereby enhance its electrochemical performance [49e51]. To assess the carbon tolerance of PBMFO anode, a constant discharge current was fixed at 0.5 A cm2 at 750  C using H2 and CH4 as fuels (Fig. 7c). After stabilizing the cell under humidified (3% H2O) H2 for 80 h, the atmosphere was changed to humidified CH4. Although the surface morphology of the particles changed slightly after operation in CH4 for 100 h due to the long thermal history, no carbon deposition was observed (Fig. S9). Importantly, the cell performance exhibited no observable degradation during operation, indicating that the layered PBMFO anode offered high carbon coking tolerance. Fig. 8 shows the output power densities and polarization resistance of the single cell using PBMFO symmetrical electrodes at 750  C after each redox cycling test, in order to evaluate the redox properties of PBMFO electrodes. During gas switch process, Ar was used as protecting gas. After undergoing fourteen redox cycles at 750  C, the maximum power output of the symmetrical SOFC remained stable, indicating that PBMFO electrodes had good redox

Fig. 8. Output power densities and polarization resistance of the PBMFOjLSGM jPBMFO single cell at 750  C after each cycling test.

L. Zhao et al. / Journal of Power Sources 342 (2017) 313e319

stability. 4. Conclusions In summary, a layered perovskite PrBaMn1.5Fe0.5O5þd (PBMFO) was studied as symmetric electrodes of SSOFCs in terms of its electrochemical property and durability. Layered PBMFO electrode demonstrated high electrical conductivity, excellent electrocatalytic activity, good redox and high carbon coking tolerance. Importantly, remarkable electrochemical performance and durability were obtained for the single cell with PBMFO symmetric electrodes using hydrocarbon and hydrogen, indicating that PBMFO could be a type of promising symmetrical electrodes for practical applications of SSOFCs. Acknowledgements The project was supported by the National Natural Science Foundation of China (Grant No. 21401171), the National Natural Science Foundation of China (Grant No. 51402266), and the Fundamental Research Funds for the Central Universities, China University of Geosciences (Beijing) (Grant No. CUG140608). Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jpowsour.2016.12.066. References [1] Y.H. Huang, R.I. Dass, Z.L. Xing, J.B. Goodenough, Science 312 (2006) 254e257. [2] S.W. Tao, J.T.S. Irvine, Nat. Mater. 2 (2003) 320e323. [3] A. Atkinson, S. Barnett, R.J. Gorte, J.T.S. Irvine, A.J. McEvoy, M. Mogensen, S.C. Singhal, J. Vohs, Nat. Mater. 3 (2004) 17e27. [4] Z.P. Shao, S.M. Haile, Nature 431 (2004) 170e173. [5] L. Yang, S.Z. Wang, K. Blinn, M.F. Liu, Z. Liu, Z. Cheng, M.L. Liu, Science 326 (2009) 126e129. [6] L. Yang, Y. Choi, W.T. Qin, H.Y. Chen, K. Blinn, M.F. Liu, P. Liu, J.M. Bai, T.A. Tyson, M.L. Liu, Nat. Commun. 2 (2011). [7] J.C. Ruiz-Morales, D. Marrero-Lopez, J. Canales-Vazquez, J.T.S. Irvine, RSC Adv. 1 (2011) 1403e1414. [8] Q.A. Liu, X.H. Dong, G.L. Xiao, F. Zhao, F.L. Chen, Adv. Mater. 22 (2010) 5478e5482. [9] S.P. Jiang, L.Z. Ab, Y. Zhang, J. Mater. Chem. 17 (2007) 2627e2635. [10] J.C. Ruiz-Morales, J. Canales-Vazquez, J. Pena-Martinez, D. Marrero-Lopez, P. Nunez, Electrochim. Acta 52 (2006) 278e284. [11] D.M. Bastidas, S.W. Tao, J.T.S. Irvine, J. Mater. Chem. 16 (2006) 1603e1605. [12] F. Napolitano, A.L. Soldati, J. Geck, D.G. Lamas, A. Serquis, Int. J. Hydrogen Energy 38 (2013) 8965e8973. [13] Y. Zheng, C. Zhang, R. Ran, R. Cai, Z. Shao, D. Farrusseng, Acta Mater. 57 (2009) 1165e1175. [14] Z. Yang, N. Xu, M. Han, F. Chen, Int. J. Hydrogen Energy 39 (2014) 7402e7406. [15] L. dos Santos-Gomez, J.M. Compana, S. Bruque, E.R. Losilla, D. Marrero-Lopez, J. Power Sources 279 (2015) 419e427.

319

[16] Q. Zhou, C. Yuan, D. Han, T. Luo, J.L. Li, Z.L. Zhan, Electrochim. Acta 133 (2014) 453e458. [17] J. Shen, Y.B. Chen, G.M. Yang, W. Zhou, M.O. Tade, Z.P. Shao, J. Power Sources 306 (2016) 92e99. [18] 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 (2015) 205e209. [19] T.H. Shin, J.H. Myung, M. Verbraeken, G. Kim, J.T.S. Irvine, Faraday Discuss. 182 (2015) 227e239. [20] T. Broux, M. Bahout, J.M. Hanlon, O. Hernandez, S. Paofai, A. Berenov, S.J. Skinner, J. Mater. Chem. A 2 (2014) 17015e17023. [21] M. Bahout, S.S. Pramana, J.M. Hanlon, V. Dorcet, R.I. Smith, S. Paofai, S.J. Skinner, J. Mater. Chem. A 3 (2015) 15420e15431. [22] F. Tonus, M. Bahout, V. Dorcet, G.H. Gauthier, S. Paofai, R.I. Smith, S.J. Skinner, J. Mater. Chem. A 4 (2016) 11635e11647. [23] S. Sengodan, S. Ahn, J. Shin, G. Kim, Solid State Ionics 228 (2012) 25e31. [24] S. Yoo, A. Jun, Y.W. Ju, D. Odkhuu, J. Hyodo, H.Y. Jeong, N. Park, J. Shin, T. Ishihara, G. Kim, Angew. Chem. Int. Ed. 53 (2014) 13064e13067. [25] L. Zhao, J.C. Shen, B.B. He, F.L. Chen, C.R. Xia, Int. J. Hydrogen Energy 36 (2011) 3658e3665. [26] S. Choi, S. Park, J. Shin, G. Kim, J. Mater. Chem. A 3 (2015) 6088e6095. [27] Y.F. Sun, Y.Q. Zhang, B. Hua, Y. Behnamian, J. Li, S.H. Cui, J.H. Li, J.L. Luo, J. Power Sources 301 (2016) 237e241. [28] 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) 1747e1753. [29] B.B. He, Z.B. Wang, L. Zhao, X. Pan, X.J. Wu, C.R. Xia, J. Power Sources 241 (2013) 627e633. [30] N. Danilovic, A. Vincent, J.L. Luo, K.T. Chuang, R. Hui, A.R. Sanger, Chem. Mater. 22 (2010) 957e965. [31] Q. Zhang, T. Wei, Y.H. Huang, J. Power Sources 198 (2012) 59e65. [32] T. Motohashi, T. Ueda, Y. Masubuchi, M. Takiguchi, T. Setoyama, K. Oshima, S. Kikkawa, Chem. Mater. 22 (2010) 3192e3196. [33] H.P. Ding, Z.T. Tao, S. Liu, J.J. Zhang, Sci. Rep. 5 (2015). [34] Y.L. Zhu, J. Sunarso, W. Zhou, Z.P. Shao, Appl. Catal. B Environ. 172 (2015) 52e57. [35] T.H. Shin, S. Ida, T. Ishihara, J. Am. Chem. Soc. 133 (2011) 19399e19407. [36] W. Zhou, R. Ran, Z. Shao, H. Gu, W. Jin, N. Xu, J. Power Sources 174 (2007) 237e245. [37] L. Zhao, B.B. He, B. Lin, H.P. Ding, S.L. Wang, Y.H. Ling, R.R. Peng, G.Y. Meng, X.Q. Liu, J. Power Sources 194 (2009) 835e837. [38] B.B. He, G. Li, L.Q. Gui, Y.H. Ling, L. Zhao, J. Mater. Chem. A 4 (2016) 4003e4008. [39] N. Trofimenko, H. Ullmann, Solid State Ionics 124 (1999) 263e270. [40] J. Van Herle, R. Ihringer, R.V. Cavieres, L. Constantin, O. Bucheli, J. Eur. Ceram. Soc. 21 (2001) 1855e1859. [41] L. Zhao, B.B. He, X.Z. Zhang, R.R. Peng, G.Y. Meng, X.Q. Liu, J. Power Sources 195 (2010) 1859e1861. [42] B.B. He, L. Zhang, Y.X. Zhang, D. Ding, J.M. Xu, Y.H. Ling, L. Zhao, J. Power Sources 287 (2015) 170e176. [43] J. Rodriguezcarvajal, Phys. B 192 (1993) 55e69. [44] Z. Jiang, Z. Lei, B. Ding, C. Xia, F. Zhao, F. Chen, Int. J. Hydrogen Energy 35 (2010) 8322e8330. [45] A. Esquirol, N.P. Brandon, J.A. Kilner, M. Mogensen, J. Electrochem. Soc. 151 (2004) A1847eA1855. [46] F. Tietz, Ionics 5 (1999) 129e139. [47] D. Lee, J.H. Han, Y. Chun, R.H. Song, D.R. Shin, J. Power Sources 166 (2007) 35e40. [48] G. Xiao, F. Chen, Electrochem. Commun. 13 (2011) 57e59. [49] W. Wang, Z. Yang, H. Wang, G. Ma, W. Gao, Z. Zhou, J. Power Sources 196 (2011) 3539e3543. [50] F. Bozza, R. Polini, E. Traversa, Electrochem. Commun. 11 (2009) 1680e1683. [51] R. Pelosato, I. Natali Sora, V. Ferrari, G. Dotelli, C.M. Mari, Solid State Ionics 175 (2004) 87e92.