LETTER. Strongly correlated perovskite fuel cells

Transcription

1 doi:1.138/nature17653 Strongly correlated perovskite fuel cells You Zhou 1, Xiaofei Guan 1, Hua Zhou 2, Koushik Ramadoss 1, Suhare Adam 1, Huajun Liu 3, Sungsik Lee 2, Jian Shi 1,4, Masaru Tsuchiya 5, Dillon D. Fong 3 & Shriram Ramanathan 1,6 Fuel cells convert chemical energy directly into electrical energy with high efficiencies and environmental benefits, as compared with traditional heat engines 1 4. Yttria-stabilized zirconia is perhaps the material with the most potential as an electrolyte in solid oxide fuel cells (SFCs), owing to its stability and nearunity ionic transference number 5. Although there exist materials with superior ionic conductivity, they are often limited by their ability to suppress electronic leakage when exposed to the reducing environment at the fuel interface. Such electronic leakage reduces fuel cell power output and the associated chemo-mechanical stresses can also lead to catastrophic fracture of electrolyte membranes 6. Here we depart from traditional electrolyte design that relies on cation substitution to sustain ionic conduction. Instead, we use a perovskite nickelate as an electrolyte with high initial ionic and electronic conductivity. Since many such oxides are also correlated electron systems, we can suppress the electronic conduction through a filling-controlled Mott transition induced by spontaneous hydrogen incorporation. Using such a nickelate as the electrolyte in free-standing membrane geometry, we demonstrate a low-temperature micro-fabricated SFC with high performance. The ionic conductivity of the nickelate perovskite is comparable to the best-performing solid electrolytes in the same temperature range, with a very low activation energy. The results present a design strategy for high-performance materials exhibiting emergent properties arising from strong electron correlations. SmNi 3 (SN) belongs to a series of rare-earth nickelates (RNi 3 or RN) with the perovskite structure (AB 3 ), which exhibits linked corner-shared B 6 octahedra (Fig. 1a) 7. In perovskite oxides, protons can form ionic defects ( H in Kröger Vink notation) by bonding with oxygen 8, and diffuse through a Grotthuss mechanism that involves the fast rotational diffusion of the protonic defects and the rate-limiting proton transfer to the neighbouring oxygen ions 8,9. The transition states of the proton rotation and proton transfer require local lattice distortions such as elongation and bending of the B bond, respectively 1,11. The schematic of proton incorporation and diffusion processes for a cubic perovskite is shown in Fig. 1b with the following processes: (i) proton incorporation, (ii) rotational diffusion, (iii) transfer to neighbouring oxygen, (iv) bending and (v) elongation of the B bond. In SN, proton incorporation and diffusion happen in a similar way, albeit with several different characteristics, as will be discussed in more detail later (Fig. 1c). In the low-temperature fuel cell operation range (3 5 C), stoichiometric SN shows metallic conductivity with an electrical resistivity of ~1 mω cm, which is detrimental to electrolyte applications. The high electronic conductivity is due to single electron occupancy on the fourfold degenerate e g manifold (including spin) on Ni 3+, as shown in Fig. 1d (in the ionic limit; the covalent limit cases are shown in Extended Data Fig. 1a and b), where carriers can migrate without overcoming the on-site Coulomb repulsion. When electrons are doped into SN via hydrogenation and the valence of nickel is reduced to Ni 2+ (with overall reaction + Ni H Ni2+ + H 2 2 ), however, electronic transport through the e g 2 manifold will be suppressed by the Hubbard intra-orbital electron electron Coulomb interaction U (Fig. 1e). Such filling-controlled Mott transitions enable the application of hydrogenated SN as an electrolyte, owing to its wide electronic bandgap 12, which is close to the Ni intra-orbital Coulomb repulsion and large enough to suppress electronic conductivity 13. Spontaneous incorporation of protons into SN upon hydrogen exposure without any electrical bias at low temperatures can be seen in Extended Data Fig. 2. This is unlike typical perovskite proton conductors such as yttrium-doped BaCe 3 and BaZr 3, where subvalent cations are needed as substitutional acceptors to facilitate the hydrogenincorporation process (Fig. 1b). Therefore the concentration of protons in SN may not be limited by the oxygen vacancy concentration, as commonly noted in acceptor-doped electrolytes. The electronic transport mechanism in H-SN is characterized by the Efros Shklovskii variable range hopping mechanism, in which small polarons form because of strong electron lattice coupling in the presence of a Coulomb gap (Extended Data Fig. 1c e). Figure 1f illustrates how this collective quantum mechanical effect enables the electrolyte design. Initially no power output is extracted from the SN-electrolyte fuel cell because of the high electronic conductivity in pristine SN. When the hydrogen fuel is introduced at the anode (catalytic Pt or Pd), hydrogen molecules dissociate into protons and donate electrons to Ni(iii) in SN at the triple phase boundaries. The hydrogenation process creates an electrically insulating H-SN on the anode side. nce this insulating layer is formed, as long as hydrogen fuel is supplied, protons can continue to diffuse under the chemical potential gradient, while the electron transport through H-SN directly to the cathode is strongly suppressed by carrier localization. As a result, electrons are forced to pass through the external circuit and generate electrical power. The time evolution of the open-circuit voltage (CV) in a micro-fabricated SFC with a free-standing SN membrane (see Extended Data Figs 3 and 4 for the device structure and fabrication) as the electrolyte verifies the above mechanism (Extended Data Fig. 5a). Initially there is no CV as the cell is electrically shorted by pristine SN. The CV increases under continuous hydrogen flow after the temperature becomes stabilized, as the H-SN phase forms on the anode side, and reaches a stable output when the stationary state is reached. The current voltage characteristics of the microfabricated SFCs (Fig. 2a) exhibit typical activation polarization, ohmic loss and concentration polarization behaviour, and the power output reaches a maximum value of 225 mw cm 2 at 5 C, which is comparable to the best-performing proton conducting fuel cells (ref. 14 and references therein). The highest CV achieved (1.3 V) is close to the Nernst potential (~1.7 V), showing that the ionic transference number is close to unity, with the electronic conduction 1 John A. Paulson School of Engineering and Applied Sciences, Harvard University, Cambridge, Massachusetts 2138, USA. 2 X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, Argonne, Illinois 6439, USA. 3 Materials Science Division, Argonne National Laboratory, Argonne, Illinois 6439, USA. 4 Department of Materials Science and Engineering, Rensselaer Polytechnic Institute, Troy, New York 1218, USA. 5 SiEnergy Systems, Cambridge, Massachusetts 214, USA. 6 School of Materials Engineering, Purdue University, West Lafayette, Indiana 4797, USA. 9 JUNE 216 VL 534 NATURE 231

2 RESEARCH LETTER a d f Sm Ni Anode Cathode Electron itinerant e g t 2g H 2 Electrolyte 2 Ni 3+ SN b (iv) H-SN SN Electron doping Mott transition Electron localized Figure 1 Solid electrolyte design principle based on the emergent phase arising from strong correlations. a, The distorted perovskite structure of the SN crystal. b, c, Proton incorporation and conduction mechanisms in a conventional solid-state electrolyte (A, B and M are metal cations and is the oxygen anion) (b) and the proposed new electrolyte (c). (i) Proton incorporation. (ii) and (iii), Proton transport by rotational diffusion within an octahedron (ii) and transfer to a neighbouring oxygen ion facilitated by the hydrogen bond (dashed red line) (iii). (iv) and (v), The bending (iv) and the stretching (v) of the metal oxygen bond promote processes (iii) and (ii), respectively. In conventional electrolytes, substitutional sub-valent cations, M, are needed to facilitate the hydrogen incorporation. In SN, proton incorporation can happen spontaneously. The ligand holes in SN reduce the effective charge of the oxygen ions (only two of them are explicitly shown). d, e, The electronic configuration of Ni 3d orbitals for the pristine (d) and the electron-doped (e) SN in the ionic limit. Electronic transport is suppressed by the on-site electron electron correlation U upon electron doping (e). f, A schematic of a SN-electrolyte SFC and its operation mechanism. Spontaneous hydrogen incorporation creates a strongly correlated insulating layer and suppresses the electronic current. TPB, triple phase boundary. almost completely suppressed. The deviation of CV from the Nernst potential in general could be related to gas leakage and residual electronic conductivity. The ionic transference number of H-SN at 5 C is estimated to be.96 using the standard electromotive force method. Because there is a small yet finite current inside the fuel cells through the electrolyte under the CV condition, the electrode polarization loss (in addition to ohmic loss) may contribute to the deviation of the measured CV from the ideal value. Therefore, a method that considers the polarization loss may also be used to evaluate the ionic transference number 15,16 (Extended Data Fig. 5b). Increasing the electrolyte thickness typically enhances the measured CV, possibly owing to the reduced possibilities of pinholes in the membrane and a decrease in the relative ratio between electrode polarization and electrolyte resistance (see Extended B 2 B Ni 3+ (2 δ) Ni 3+ (iii) H + (v) (iii) H + A (ii) 2 (iv) Sm (ii) (2 δ) (v) B V M (i) Ni 3+ (i) Ni 2+ H 2 H 2 TPB e c U e H 2 Ni 2+ H-SN e e g t 2g Pt H H + e Data Fig. 6a and additional discussions in the Supplementary Information). H-SN fuel cells with dense Pd anodes also produce power output, indicating that protons rather than oxygen ions are the dominant mobile ion species in the material (Extended Data Fig. 6b). In addition, these fuel cells can also work under pure H 2 and are stable over tens of hours (Extended Data Figs 6b and 7). The Nyquist plot of the cell under CV conditions measured at 5 C is shown in Fig. 2b. The extrapolated area specific resistance of H-SN, one of the key performance metrics, is remarkably low:.45 Ω cm 2 at 5 C, which is less than one-third of the general target value (.15 Ω cm 2 ) for the area specific resistance of an oxide electrolyte 17. The Nyquist plot can be modelled by an equivalent circuit with an ohmic resistor, R hm, and serial elements each consisting of a resistor, R i (i = 1, 2, 3), and a constant phase element, CPE i (i = 1, 2, 3), as summarized in Supplementary Table 1. The three semicircles in the Nyquist plot originate from the electric double-layer capacitance at the anode and cathode, and the pseudocapacitance related to hydrogen incorporation in SN (see the footnote to Supplementary Table 1). Figure 2c shows the ionic conductivity (calculated by the electromotive force method) of H-SN measured from free-standing Pt/H-SN/Pt micro-fabricated SFCs and H-SN epitaxial films on LaAl 3 (LA) (1) (indexed in pseudocubic notation), compared with several other best-performing oxygen-ion-conducting and proton-conducting electrolyte materials SN has a high ionic conductivity with low activation energy (~.3 ev, similar to solid acid protonic conductors 22 ), making it especially suitable for lowtemperature SFC applications 3. The difference in the ionic conductivity measured from epitaxial thin films and membranes could be related to contributions from grain boundaries 2,23. Grain boundaries may not only decrease proton mobility by scattering and trapping, but may also reduce the proton concentration proximal to the boundaries by creating space charge layers. Therefore the total ionic resistance of polycrystalline samples can be larger than that of the epitaxial films. Several factors may collectively lead to the high ionic conductivity with low activation energy in SN. First, it has been found that in RN, Ni forms a covalent bond with in a mixed electron configuration of 3d 7 and 3d 8 L (where L denotes a ligand hole on 2p) (ref. 24). The covalence reduces the effective charge on oxygen and therefore the bonding strength between the oxygen ion and the proton, which lowers the proton transfer activation energy (Fig. 1c). Additionally, the proton transport barrier in perovskites with a tetravalent B-site (A(ii) B(iv)) is in general much smaller than the ones with a pentavalent B-site (A(i) B(v)) 25. It has been suggested that A(iii) B(iii) perovskites may have even higher ionic conductivity 25. This may be explained by the weaker repulsion between B-site ions and protons in A(iii) B(iii) perovskites, which reduces the energy of the proton in its transition state. Finally, as the transition states of the proton rotation and proton transfer require local lattice distortions such as elongation and bending of the B bond, respectively 1,11, the relative low energy of the Ni bending and stretching modes in SN (~35 mev and 75 mev, respectively 26 ) can also contribute to lowering the proton transport barrier. To confirm the electron localization mechanism during fuel cell operation and to reveal the underlying reasons for the high ionic conductivity, both chemical and structural characterizations of the SN hydrogenation process were performed. Ex situ X-ray absorption near-edge spectroscopy (XANES) measurements of the nickel K-edge from a pristine and a hydrogenated SN sample are shown in Fig. 3a, as well as that from a reference nickel metal sample used for energy calibration. Several features are present in the spectra of SN and H-SN. The pre-edge feature, A, originates from the dipolar transition between Ni 1s and Ni 3d 2p hybridized 3d 8 L, and points to the covalent nature of the Ni bond 27. Features B, D and E are derived from the first oxygen coordination shell, while C and C originate from the second shell of the rare-earth ions NATURE VL JUNE 216

3 RESEARCH a Voltage (V) Cell 3 Cell 1 Cell ,2 1,6 Current density (ma cm 2 ) Power density (mw cm 2 ) b Im(Z) (Ω cm 2 ) Hz Figure 2 Performance of the emergent-phase electrolyte in fuel cells. a, Typical current density voltage characteristics and the power densities of Pt/H-SN/Pt micro-fabricated SFCs measured at 5 C with 3% humidified 5% H 2 95% Ar as fuel and laboratory air as oxidant. The electrolyte thickness is 1.5 μm for cell 1, and 1 μm for cells 2 and 3. b, A Nyquist plot measured under CV conditions at 5 C for a Pt/SN/ Pt cell (solid line shows the fitted curve). Z is the complex impedance measured from the fuel cell. The Nyquist plot can be modelled by an equivalent circuit (inset) with an ohmic resistor, R hm, and serial R hm R 1 R 2 R 3 CPE 1 CPE 2 CPE Hz 1 Hz Re(Z) (Ω cm 2 ) c log (S cm 1 ) Temperature (ºC) Epitaxial H-SN Highly textured BZY film BCY 1.2 BZY YSZ ,/T (K 1 ) Pt/H-SN/Pt LSGM GDC elements each consisting of a resistor, R i (i = 1, 2, 3), and a constant phase element, CPE i (i = 1, 2, 3). c, The ionic conductivity of H-SN compared to the best-performing oxygen-ion-conducting electrolytes (dashed lines) and proton conductors (solid lines). The oxygen-ionconducting electrolytes are: stabilized zirconia (YSZ, (Zr 2 ).9 (Y 2 3 ).1 ) (ref. 17), La.8 Sr.2 Ga.8 Mg.2 3 (LSGM) (ref. 18) and doped ceria (GDC, Ce.8 Gd δ ) (ref. 19). The proton conductors are BaZr.8 Y.2 3 δ (BZY, in the form of both sintered pellets and highly textured films) (ref. 2) and BaCe.8 x Zr x Y.2 3 δ (BCY, < x <.8) (ref. 21). A substantial shift of the absorption edge to a lower energy is observed upon hydrogenation. The energies of the absorption edge and other features are consistent with those of Ni(iii) in RN 28. From the inflection point in the first derivative of the absorption (Fig. 3b), the chemical shift from SN to H-SN is determined to be ~2. ev. A linear relation between the absorption edge and the formal valence state has been previously noted with a slope from ~1.5 ev per electron to ~2.8 ev per electron The absorption inflection point, however, depends on both the Ni valence state and the atomic arrangement around Ni. In this study, the valence change is the primary factor leading to the absorption edge shift, and the Ni valence state change is close to 1, suggesting high proton concentration in H-SN without introducing impurity dopants. The change of Ni valence state verifies that hydrogen exists as protons in H-SN, because it is more favourable for Ni to accept an electron from hydrogen rather than from 2 and Sm 3+ ions when changing from SmNi(iii) 3 and H-SmNi(ii) 3. The angle-resolved XANES spectra show that the proton incorporation not only happens at the surface but also through the thickness of the films (Extended Data Fig. 8). In addition, the decrease in the white line intensity suggests an overall decrease in the hole density on Ni after hydrogenation. The intensity of the pre-edge feature A, offset with respect to feature B by ~14.4 ev, represents the density of the 3d 8 L state in the ground state and decreases upon hydrogenation (inset of Fig. 3a), which shows that the doped electrons partially fill the ligand holes. Following the analysis in ref. 27, we find the pristine ground state to be ~.5 3d d 8 L, and estimate that the concentration of ligand holes decreases by ~5% after the hydrogenation. This verifies that ligand holes are present on oxygen ions in both SN and H-SN, which helps to reduce the proton transfer activation energy. Figure 3c shows the representative in situ XANES spectra during the hydrogenation process. The chemical shift is smaller than those of the ex situ experiments owing to the lower operation temperature limited by the apparatus. The dynamic change in the absorption edge (inset of Fig. 3c) shows that the average valence state reaches equilibrium in ~3 min at 2 C. Synchrotron X-ray diffraction studies (Extended Data Figs 9 and 1) suggest that the SN lattice expands during hydrogenation, which may lead to a change in the relative rate of inter- and intra-octahedron proton transfer 9 and the lattice open volume, and therefore modify the long-range proton transport properties. a Normalized absorption (arbitrary units) Ni foil H-SN SN A B.1 C C A 8,32 8,34 8,36 8,38 8,4 Photon energy (ev) Figure 3 Ex situ and in situ XANES characterizations of the phase evolution. a, Ex situ normalized Ni K-edge XANES spectra of SN, H-SN and the nickel metal reference, with zoomed view of the pre-edge feature A (inset). The other features (B, C, C, D and E) are derived from the first oxygen coordination shell and the second shell of the rare-earth D 8,33 8,335 E b First derivative (arbitrary units) Ni foil H-SN SN 8,32 8,34 8,36 8,38 8,4 Photon energy (ev) ions. b, First derivative of the normalized absorption. c, In situ XANES spectra of the SN hydrogenation process performed at 2 C. The arrow indicates the direction of time evolution. The dynamics of the shift in the energy of the absorption edge E k is shown in the inset (where E k, is the absorption edge energy of pristine SN). 8,32 8,34 8,36 8,38 8,4 Photon energy (ev) c Normalized absorption (arbitrary units) In situ hydrogenation E k E k, (ev) Time (min) 9 JUNE 216 VL 534 NATURE 233

4 RESEARCH LETTER nline Content Methods, along with any additional Extended Data display items and Source Data, are available in the online version of the paper; references unique to these sections appear only in the online paper. Received 13 September 215; accepted 1 March 216. Published online 16 May Hayre, R. P., Cha, S.-W., Colella, W. & Prinz, F. B. Fuel Cell Fundamentals (John Wiley & Sons, 26). 2. Singhal, S. Advances in solid oxide fuel cell technology. Solid State Ion. 135, (2). 3. Wachsman, E. D. & Lee, K. T. Lowering the temperature of solid oxide fuel cells. Science 334, (211). 4. Shao, Z. et al. A thermally self-sustained micro solid-oxide fuel-cell stack with high power density. Nature 435, (25). 5. Haile, S. M. Fuel cell materials and components. Acta Mater. 51, (23). 6. Kerman, K., Lai, B.-K. & Ramanathan, S. Free standing oxide alloy electrolytes for low temperature thin film solid oxide fuel cells. J. Power Sources 22, (212). 7. Catalan, G. Progress in perovskite nickelate research. Phase Transit. 81, (28). 8. Kreuer, K.-D. Proton conductivity: materials and applications. Chem. Mater. 8, (1996). 9. Münch, W., Kreuer, K.-D., Seifert, G. & Maier, J. Proton diffusion in perovskites: comparison between BaCe 3, BaZr 3, SrTi 3, and CaTi 3 using quantum molecular dynamics. Solid State Ion. 136/137, (2). 1. Münch, W., Seifert, G., Kreuer, K. D. & Maier, J. A quantum molecular dynamics study of proton conduction phenomena in BaCe 3. Solid State Ion , (1996). 11. Münch, W., Kreuer, K. D., Seifertli, G. & Majer, J. A quantum molecular dynamics study of proton diffusion in SrTi 3 and CaTi 3. Solid State Ion. 125, (1999). 12. Shi, J., Zhou, Y. & Ramanathan, S. Colossal resistance switching and band gap modulation in a perovskite nickelate by electron doping. Nat. Commun. 5, 486 (214). 13. Tuller, H. L. Defect engineering: design tools for solid state electrochemical devices. Electrochim. Acta 48, (23). 14. Duan, C. et al. Readily processed protonic ceramic fuel cells with high performance at low temperatures. Science 349, (215). 15. Liu, M. & Hu, H. Effect of interfacial resistance on determination of transport properties of mixed-conducting electrolytes. J. Electrochem. Soc. 143, L19 L112 (1996). 16. Wang, S., Wu, L., Gao, J., He, Q. & Liu, M. xygen ion transference number of doped lanthanum gallate. J. Power Sources 185, (28). 17. Steele, B. C. & Heinzel, A. Materials for fuel-cell technologies. Nature 414, (21). 18. Ishihara, T., Shibayama, T., Honda, M., Nishiguchi, H. & Takita, Y. Intermediate temperature solid oxide fuel cells using LaGa 3 electrolyte II. Improvement of oxide ion conductivity and power density by doping Fe for Ga site of LaGa 3. J. Electrochem. Soc. 147, (2). 19. Esposito, V. & Traversa, E. Design of electroceramics for solid oxides fuel cell applications: playing with ceria. J. Am. Ceram. Soc. 91, (28). 2. Pergolesi, D. et al. High proton conduction in grain-boundary-free yttriumdoped barium zirconate films grown by pulsed laser deposition. Nat. Mater. 9, (21). 21. Fabbri, E., D Epifanio, A., Di Bartolomeo, E., Licoccia, S. & Traversa, E. Tailoring the chemical stability of Ba(Ce.8 x Zr x )Y.2 3 δ protonic conductors for intermediate temperature solid oxide fuel cells (IT-SFCs). Solid State Ion. 179, (28). 22. Haile, S. M., Boysen, D. A., Chisholm, C. R. I. & Merle, R. B. Solid acids as fuel cell electrolytes. Nature 41, (21). 23. Shim, J. H., Gür, T. M. & Prinz, F. B. Proton conduction in thin film yttriumdoped barium zirconate. Appl. Phys. Lett. 92, (28). 24. Medarde, M. L. Structural, magnetic and electronic properties of RNi 3 perovskites (R = rare earth). J. Phys. Condens. Matter 9, 1679 (1997). 25. Kreuer, K. n the complexity of proton conduction phenomena. Solid State Ion. 136/137, (2). 26. Jaramillo, R., Ha, S. D., Silevitch, D. M. & Ramanathan, S. rigins of bad-metal conductivity and the insulator-metal transition in the rare-earth nickelates. Nat. Phys. 1, (214). 27. García, J., Blasco, J., Proietti, M. G. & Benfatto, M. Analysis of the x-rayabsorption near-edge-structure spectra of La 1 x Nd x Ni 3 and LaNi 1 x Fe x 3 perovskites at the nickel K edge. Phys. Rev. B 52, (1995). 28. Medarde, M. et al. Charge disproportionation in RNi 3 perovskites (R = rare earth) from high-resolution x-ray absorption spectroscopy. Phys. Rev. B 8, (29). 29. Tan, Z., Heald, S. M., Cheong, S. W., Cooper, A. S. & Moodenbaugh, A. R. Nature of hole doping in Nd 2 Ni 4 and La 2 Ni 4 : Comparison with La 2 Cu 4. Phys. Rev. B 47, (1993). 3. Grady, W. E., Pandya, K. I., Swider, K. E. & Corrigan, D. A. In situ x-ray absorption near-edge structure evidence for quadrivalent nickel in nickel battery electrodes. J. Electrochem. Soc. 143, (1996). Supplementary Information is available in the online version of the paper. Acknowledgements Financial support was provided by the Army Research ffice (grants W911NF and W911NF ), the Air Force ffice of Scientific Research (grant FA ), the Advanced Research Projects Agency-Energy (ARPA-E), an IBM PhD Fellowship and the National Academy of Sciences. Part of the work was performed at the Center for Nanoscale Systems at Harvard University. Use of the Advanced Photon Source was supported by the US Department of Energy, ffice of Science, ffice of Basic Energy Sciences, under contract number DE-AC2-6CH D.D.F. was supported by the U.S. Department of Energy, ffice of Science, ffice of Basic Energy Sciences, Materials Sciences and Engineering Division. Author Contributions Y.Z. and S.R. conceived the study. Y.Z. fabricated the fuel cells and performed the initial tests. X.G. designed and performed the quantitative fuel-cell tests and analysis. Y.Z., H.Z., H.L. and S.L. performed the X-ray absorption spectroscopy measurements. Y.Z. and H.Z. conducted the X-ray diffraction characterizations. K.R. performed the low-temperature electronic transport measurements. S.A. prepared the freestanding Si 3 N 4 membrane. M.T. provided technical advice on the micro-sfc fabrication and characterization. Y.Z., X.G., J.S. and S.R. wrote the manuscript. All authors discussed the results and commented on the manuscript. Author Information Reprints and permissions information is available at The authors declare no competing financial interests. Readers are welcome to comment on the online version of the paper. Correspondence and requests for materials should be addressed to S.R. (shriram@purdue.edu) or Y.Z. (youzhou@fas.harvard.edu). 234 NATURE VL JUNE 216

5 RESEARCH METHDS Micro-fabricated SFC. 4-inch-diameter, 525-μm-thick Si (1) wafers coated with 2 nm Si 3 N 4 on both sides were used as substrates for micro-fabrication of an SFC. ne side of the wafer was patterned with photolithography to define silicon nitride areas uncovered by photoresist. Then the uncovered silicon nitride was removed by reactive ion etching in CF 4 and 2. Afterwards, the exposed Si was etched with a 3 wt% KH aqueous solution at 86 C for ~5 h to leave a μm 2 free-standing Si 3 N 4 membrane. After KH etching, the 4-inch wafers were cut into 1 1 cm 2 chips, with nine windows on each of the chips. SN electrolyte films with thicknesses ranging from 5 nm to 1.5 μm were deposited onto the silicon nitride membranes by radio-frequency magnetron sputtering in an Ar/ 2 mixture at a total pressure of 5 mtorr, either from a ceramic SN target, or from two metallic Ni and Sm targets. For films sputtered from metal targets, the chips were annealed under 1 bar of pure 2 at 5 C for 24 h so that SN would form the perovskite phase after annealing. The growth rate of SN was calibrated by X-ray reflectivity, cross-section transmission electron microscopy and scanning electron microscopy. The Sm:Ni cation ratio was determined by energy-dispersive X-ray spectroscopy. Then, the Pt cathode was deposited by magnetron sputtering in pure Ar at a total pressure of 75 mtorr, which yields a porous Pt layer to increase the size of the triple phase boundaries on the anode and cathode side. The Si 3 N 4 layer on the backside was removed by reactive ion etching in CF 4 and 2. Finally, a Pt anode layer was deposited into the Si well side using sputtering under the same conditions as for the cathode. Dense palladium films of thickness 1 2 nm were also used as the anode, which were deposited by an electron-beam evaporator. The detailed fabrication process is shown in Extended Data Fig. 3. Synthesis of SN epitaxial thin films. The epitaxial SN thin films were grown on LA (1) by radio-frequency magnetron sputtering in an Ar/ 2 mixture at total pressure of 5 mtorr from two metallic Ni and Sm targets. The samples were sealed in a vessel under 1 bar of pure 2 and annealed at 5 C for 24 h in a tube furnace. Electrical and electrochemical characterization. Fuel-cell tests were performed in a custom-design fuel cell test station. The morphology of the membranes during fuel cell testing was monitored in situ under an optical microscope. Anode current was collected with a gold -ring and a stainless-steel base, and the cathode current was collected through a micromanipulator probe with a Pt-plated tungsten tip. The electrochemical active area for fuel cell performance was defined as the area of the free-standing SN membrane. For epitaxial thin films, the conductivity measurements were done using in-plane geometry with porous Pt electrodes. The current voltage characteristics were measured by starting at an CV and sweeping down to V at a rate of 2 mv s 1 (or 1 mv s 1 ). Electrochemical impedance spectroscopy was scanned from 1 6 Hz to 1 Hz with an amplitude of 2 mv. All the electrochemical measurements were performed with a Solartron 126/1287 electrochemical test setup. The impedance data were fitted using ZView software. For Pt/SN/Pt fuel cells, either dry or moist 5% H 2 /95% Ar was flown onto the anode side. For Pt/SN/Pd fuel cells, pure H 2 bubbled through room-temperature water was flown onto the Pd anode. In both cases, stationary air was used as the cathode oxidant. The ionic conductivity of epitaxial thin films was measured in dry 5% H 2 /95% Ar and the conductivity of suspended membranes in Pt/H-SN/ Pt SFCs was measured with 3% humidified 5% H 2 /95% Ar as fuel and laboratory air as oxidant. In situ conductivity dynamics measurements were performed with a Keithley 2635A and Solartron 126/1287 in a custom-built chamber by switching between dry 5% H 2 /95% Ar and 2 with a fixed flow rate of 15 standard cubic centimetres per minute (sccm). Electronic transport studies below room temperature were performed in vacuum ex situ using a Lakeshore probe-station and a Keithley 2635A on samples annealed in dry 5% H 2 /95% Ar for 3 min at 2 C. X-ray absorption spectroscopy studies. The X-ray absorption spectroscopy data were acquired at the bending magnet beamline, 12-BM-B, at the Advanced Photon Source, Argonne National Laboratory. The absorption was measured in fluorescence mode with the samples placed in a custom-made cell allowing in situ control of the atmosphere and heating of the sample. An infrared heater is used to heat the sample up to 2 C. A 13-element Ge detector (Canberra) was used to measure the fluorescence yield. Grazing incidence geometry was used to minimize the elastic scattering intensity. The incident angle is varied from.25 to 5, covering a range below and above the critical angle. The calibration of the monochrometer was monitored by simultaneously measuring the absorption of a nickel reference foil during each measurement. For ex situ XANES measurements, SN samples annealed in dry 4% H 2 /96% Ar for 3 min at 2 C or 3 C. The data were normalized by fitting the pre-edge to zero and the post-edge to 1 using Ifeffit performed by the software Athena ( Both epitaxial SN thin films of different thickness on LA and polycrystalline SN thin films on Si 2 /Si were characterized by XANES. Synchrotron X-ray diffraction. Synchrotron X-ray diffraction of the SN samples were conducted at an insertion device beamline, 12ID-D at the Advanced Photon Source on a six-circle Huber goniometer with an X-ray energy of 2 kev using a pixel array area detector (Dectris Pilatus 1 K). The X-ray beam had a flux of 1 12 photons per second. The q z -scan (L-scan) was obtained by removing the background scattering contributions using the two-dimensional images. For ex situ X-ray diffraction measurements, SN samples were grown on LA substrates and annealed in 5% H 2 /95% Ar at 3 C for 2 h. For the real-space mapping shown in Extended Data Fig. 9d f, an X-ray footprint of 5 μm (horizontal in Extended Data Fig. 9b) 5 μm (vertical in Extended Data Fig. 9b) was used to scan across the sample, collecting the diffraction pattern from each point. SN stability test. To test the material stability in a pure hydrogen atmosphere, we annealed SN thin films under 1 bar of pure H 2 in a tube furnace at 5 C for 48 h and 72 h. Pt electrodes were deposited onto SN thin films as catalyst. The H 2 flow was set to a constant of 3 sccm. X-ray diffraction was performed on the annealed samples using a Bruker-D8 Discover diffractometer. 31. Shklovskii, B. I. & Efros, A. L. Electronic Properties of Doped Semiconductors Ch. 9/1, (Springer, 1984). 32. Goodenough, J. B. Electronic and ionic transport properties and other physical aspects of perovskites. Rep. Prog. Phys. 67, 1915 (24). 33. Natoli, C. R. in EXAFS and Near Edge Structure III Vol. 2 Springer Proceedings in Physics (eds Hodgson, K.., Hedman, B. & Penner-Hahn, J. E.) Ch. 1, (Springer, 1984).

6 RESEARCH LETTER Extended Data Figure 1 The electronic structure of SN and H-SN in the covalent limit and their electronic transport mechanisms. a, The electronic structure of SN in the covalent limit. Ligand holes are present on 2p orbitals of the pristine SN, while two electrons occupy the Ni e g manifold. The pristine SN, however, is not strongly correlated because carriers transport through the 2p ligand holes. b, The electronic structure of H-SN. Upon electron doping and thus filling of the ligand holes, electrons have to overcome Hubbard intra-orbital correlation U to transport, which opens up a large Mott gap, and suppresses the electronic conduction in SN. c, The resistivity ρ of H-SN compared with pristine SN. The resistivity of H-SN is more than eight orders of magnitude larger than that of pristine SN at room temperature. d, e, Derivatives of resistivity ( dlnρ/dlnt) as a function of T plotted in log log scale for H-SN and SN. The transport mechanism can be determined from the slope p of the dlnρ/dlnt versus T curves. H-SN shows the Efros Shklovskii variable range hopping mechanism (p = 1/2), indicating polaron formation in the presence of a Coulomb gap 31 (d). Pristine SN shows crossover from activated conduction (p = 1) to Mott variable range hopping (p = 1/4) (e). The Coulomb repulsion is less strong in pristine SN.

7 RESEARCH Extended Data Figure 2 Fuel-induced suppression of electronic conduction in SN. a, Temporal evolution of SN conductivity when switching between different gas environments at various temperatures. b d, Images of SN and H-SN on transparent substrate LA. b, Pristine SN shows dark, shining colour and the Pt bars are bright. c, After annealing in 5% H 2 /95% Ar at 3 C for 1.5 h and cooling down to room temperature in the same gas environment, SN near the Pt electrodes becomes electronically insulating and transparent. A clear diffusion profile can be seen as the transparent region has a shape similar to the outline of the Pt electrodes. d, An optical micrograph of the hydrogenated SN indicates a diffusion profile of protons from the triple phase boundaries. The diffusion length L D is estimated to be ~3 μm.

8 RESEARCH LETTER Extended Data Figure 3 A schematic of the fabrication process of fuel cells with free-standing SN membranes as the electrolyte. a, Patterning a etch mask on the back side of the Si 3 N 4 /Si/Si 3 N 4 chip with photoresist (PR). b, Removing exposed silicon nitride by reactive ion etching in CF 4 and 2. c, Etching the Si from the back side with a KH aqueous solution to make free-standing Si 3 N 4 membrane. d, Depositing SN thin films onto the Si 3 N 4 membranes by radio-frequency magnetron sputtering and post-annealing the sample to form stoichiometric SN. e, Fabricating the porous Pt cathodes on the front side of the chip. f, Removing the silicon nitride membrane from the back side of the chip to expose SN, using reactive ion etching. g, h, Depositing anodes on the back side of the chip. Two types of fuel cell anodes were studied in this work: porous Pt as a model system (g) and a dense Pd anode (h). Pd is an industry-standard proton conducting membrane that is used in this study to selectively permeate protons from the fuel side to the cathode.

9 RESEARCH Extended Data Figure 4 SN micro-fabricated SFCs and fuel cell test apparatus. a, An image of a 1 mm 1 mm Si 3 N 4 /Si chip with nine SN-electrolyte fuel cells (US dime coin shown for size). b, c, ptical micrographs of the free-standing buckled SN membrane due to local compressive strain with top Pt cathode on a Si chip. The buckled morphology is due to local compressive strain, engineered intentionally by synthesis and is critical for the mechanical stability and performance of the SFC. d, A scanning electron microscope of the top porous Pt electrode. e, A schematic of the customized low-temperature micro-fabricated SFC (μsfc) testing station. Both pure H 2 and 5%H 2 /95% Ar were used as fuel in the experiments.

10 RESEARCH LETTER Extended Data Figure 5 CV of H-SN micro-fabricated SFCs. a, Temporal evolution of the CV of a Pt/H-SN/Pt micro-fabricated SFC with a 3% humidified 5% H 2 /95% Ar as fuel and laboratory air as oxidant, as the temperature ramps up. Initially SN is electronically conductive so the CV is close to zero. During the hydrogenation process, the CV continues to increase after the temperature is stabilized and reaches near-ideal CV, indicating that electronic conduction is almost completely suppressed in H-SN by the Mott transition. The hydrogen fuel was always supplied at a constant flow rate both before t = and during the experiments, and the initial low CV is not due to the lack of fuel. b, The ionic transference number of H-SN at various temperatures of interest to low-temperature SFCs measured in Pt/SN/Pt cells. Two methods can be used to calculate the ionic transference number. In the electromotive force (E.M.F.) method, the fuel cell under the CV condition (infinitely large external resistance load) is modelled with an equivalent circuit containing a voltage source with an output voltage of Nernst potential E N, and two resistors R ion and R e, which correspond to the electrolyte s ionic resistance and electronic resistance, respectively. The equivalent circuit is similar to the one shown in the inset, but without the R polarization element (drawn in red). V C is the measured CV. Note that there will be a small leakage current i leak due to the finite electronic resistance of the electrolyte, but the electromotive force method assumes that the interface processes are infinitely fast and omits the polarization loss. In the method developed by Liu et al. 15, since there is a very small leakage electronic current flowing through the electrolyte, one needs to consider the electrode polarization loss. Therefore, an extra resistive element (R polarization ) needs to be considered in the equivalent circuit as shown in the inset (the red-coloured element corresponds to the extra term). With reduced polarization and increased electrolyte resistance, the ionic transference number calculated by the two methods tends to converge (see Supplementary Information for more discussion).

11 RESEARCH Extended Data Figure 6 H-SN fuel cell performance. a, The dependence of micro-fabricated SFC performance on the thickness of the SN electrolyte at 5 C. We fabricated a series of samples with various thicknesses of the electrolyte while keeping identical deposition conditions for the cathode and anode. By doing so, the electrolyte hmic resistance is varied while the electrode polarization resistance is kept more or less a constant. A clear increase in CV with increasing thickness can be seen, which could be due to the decrease in the electrode polarization loss because of the larger electrolyte hmic resistance, as discussed in Extended Data Fig. 5. The power density does not show much dependence on the electrolyte thickness, because thicker electrolytes leads to higher hmic resistance, but also higher CV. b, Performance of Pt/SN/Pd micro-fabricated SFCs with a dense Pd anode with 3% humidified pure H 2 as fuel and laboratory air as oxidant. It has been shown that hydrogen primarily creates protonic defects rather than oxygen vacancies in SN (ref. 12). To verify that protons are the dominant mobile ion species in SN and H-SN, we fabricated an SFC with the SN electrolyte, a dense 1-nm-thick Pd anode, and a porous 1-nm-thick Pt cathode. Pd anode is known as a protonic conductor but an oxygen ion barrier and can therefore filter out any oxygen ion transport. This verifies that protons rather than oxygen ions are the dominant mobile ions in SN. During the fuel cell testing, 1 sccm pure H 2 was flowed on the anode side, with the cathode exposed to air. The fuel cell with dense Pd has an CV of.6 V and a peak power density of 24 mw cm 2 at 5 C. The protonic conductivity of H-SN can be extrapolated from impedance spectroscopy and CV measurements. The similar values of the measured ionic conductivity in cells with Pt and Pd anode confirm that protonic conduction is the dominant ionic transport mechanism.

12 RESEARCH LETTER Extended Data Figure 7 Stability of H-SN. a, Cell voltage measured at 5 C for a Pt/SN/Pd fuel cell with wet 1% H 2 as the fuel and stationary air as oxidant with current being ma cm 2 (CV condition) and 78 ma cm 2, respectively. The operation is stable for more than 2 h, implying that H-SN exhibits considerable stability for fuel cell operation. The power output decreases slightly as a function of time owing to coarsening-induced porosity reduction of the metallic electrodes when current is drawn at 5 C. b, X-ray diffraction pattern of SN, and H-SN (on LA substrates) after being annealed under 1 bar of pure H 2 at 5 C for 48 h and 72 h. No new diffraction peaks are observed after annealing, which shows that H-SN is quite stable in pure H 2 for extended periods of time. θ is the incident angle of the X-ray.

13 RESEARCH Extended Data Figure 8 Angle-dependent XANES characterization. a, Ex situ angle-dependent XANES spectra of hydrogenated SN with a reference spectrum from pristine SN. The critical angle θ c of X-ray scattering for SN at the X-ray energy near Ni K-edge is calculated to be.335. When the X-ray incident angle is below the critical angle (.25 ), the XANES signal is surface sensitive with a penetration depth of ~1 nm. For an incident angle of 5, the penetration depth is close to 1 μm. The absence of angle-dependence of the XANES spectra shows that the hydrogen incorporation happens almost homogeneously across the film thickness. The XANES spectrum acquired at incident angle of 1 (not shown) is also similar to those at.25 and 5. b, The first derivative of the normalized absorption shows a similar change in the average valence state of Ni at the film surface and in the bulk.

14 RESEARCH LETTER Extended Data Figure 9 Synchrotron structural characterization of the emergent SN phase. a, An increase in the lattice constant can be caused by the larger crystal radius of Ni 2+ and electron localization. When the formal valence state of Ni reduces, its ionic radius R Ni increases, leading to the elongation of the Ni bond. In addition to the simple valencestate-related lattice expansion, electron localization can also increase the metal oxygen bond length, which can be understood on the basis of the virial theorem for central-force fields: 2 T + V =, where T is the mean kinetic energy of electrons, and V is the average potential energy. When transiting from itinerant to localized electronic behaviour, the absolute value V must decrease, which is achieved by a longer metal oxygen bond length 32, that is, Ni loc exceeds Ni itin even for the same valence state. b, An optical image of a hydrogenated SN sample. H-SN phase forms near and under the Pt electrodes, while a part of the sample remains in its pristine phase. c, X-ray diffraction patterns from the various spots A, B, C and D marked in b. The SN and LA peaks are indexed in pseudocubic notation. As the pristine SN has a pseudocubic lattice constant close to that of the LA, the SN (2) appears almost as a shoulder of the LA (2) peak. With decreasing distance between the X-ray spot and Pt electrodes, SN (2) indeed shifts to smaller q z (no other peaks observed). Two peaks (peak 1 at q z = 3.18 Å 1 and peak 2 at q z = 2.98 Å 1 ) appear in the hydrogenated region and correspond to ~4% and ~1% increase in the lattice constant. Peak 2 has the largest intensity right underneath the Pt catalyst, while peak 1 has the highest intensity far away from the Pt electrodes. The difference in the lattice constant change can be related to the decreasing doping concentration with increasing diffusion length from the triple phase boundary where hydrogen enters SN (Extended Data Fig. 1c). d f, Real-space mapping of the intensity of the Pt (111) peak at q z = 2.78 Å 1 (d), the H-SN peak 1 at q z = 3.18 Å 1 (e) and the H-SN peak 2 at q z = 2.98 Å 1 (f). A clear positive correlation between the Pt (111) and the q z = 2.98 Å 1 peaks can be seen, whereas the Pt (111) and q z = 3.18 Å 1 peaks show a negative correlation. The intensity of both peaks 1 and 2 is low in the pristine region, as expected. The increase in the average Ni bond length can be also inferred from XANES spectra using Natoli s rule 33, which states that the energy separation between features B, D, and E will scale inversely with the square of the Ni distance, because they are derived from the first oxygen coordination cell 27.

15 RESEARCH Extended Data Figure 1 Raw X-ray diffraction patterns and a schematic of proton diffusion. a, b, The collected raw two-dimensional diffraction patterns for the real-space mapping in Extended Data Fig. 9. To get the real-space mapping of different peaks across the sample, we scan the sample with an X-ray footprint of 5 μm (horizontal in Extended Data Fig. 9b) 5 μm (vertical in Extended Data Fig. 9b) by collecting the diffraction pattern from each point with an area detector. Then we calculate the diffraction intensity of each peak (Pt (111), and peak 1, 2 in Extended Data Fig. 9c) at each real-space spot from the 2d images and map it into real space to create Extended Data Fig. 9d f. a, Diffraction pattern of Pt (111) from a spot on the Pt electrode. A diffraction ring is observed as Pt is polycrystalline. b, Diffraction pattern at q z = 3.18 Å 1 from a spot between the Pt electrodes. Unlike the Pt pattern, it shows up as a point with a truncation rod rather than a ring in k-space, indicating that H-SN is still epitaxial on LA after hydrogenation. For both a and b the region inside the white dashed line was used to calculate the signal, while the region enclosed by the red dashed line but not by the white dashed line was used to calculate the background along both the q z and q x directions. The signal/background region and calculation algorithm were kept the same for all the real-space spots measured on the sample for a particular spot in the reciprocal space. c, A schematic of proton incorporation and diffusion near Pt electrodes. The part of SN directly underneath the porous Pt electrodes is on average closer to the triple phase boundaries (TPBs) than the SN region between the Pt electrodes. Therefore, a higher concentration of protons is expected under the Pt electrodes, which explains the larger lattice constant change and the correlation relation found in Extended Data Fig. 9. As the thickness of the film (z ~ 1 nm) is much smaller than the diffusion length (hundreds of micrometres), the proton concentration should not vary much along the thickness direction for the case of epitaxial thin films on LA.