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1 The effect of mechanical twisting on oxygen ionic transport in solid-state energy conversion membranes by Yanuo Shi 1, Alexander Hansen Bork 1, Sebastian Schweiger 1, Jennifer Lilia Marguerite Rupp 1 1 ) Electrochemical Materials, Department of Materials, ETH Zurich, Switzerland Correspondence should be addressed to J.L.M.R. S1: In-situ heating x-ray diffraction analysis of Gd0.2Ce0.8O1.9-x thin film For material characterization and to probe thermal expansion, in-situ heated X-ray diffraction is used to analyze the Gd0.2Ce0.8O1.9-x thin films on Si 3 N 4 /Si substrates. In Figure S1, the θ-2θ diffraction scans of a Gd0.2Ce0.8O1.9-x thin film are displayed heated between 20 and 600 C at 5 C/min. The characteristic Gd 0.2 Ce 0.8 O 1.9-x diffraction peaks of (111), (200), (311), (222) indicate that the film is polycrystalline and of cubic fluorite structure. This is in agreement with literature on pulsed laser deposited Gd0.2Ce0.8O1.9-x film at room temperature 1. With an increase in temperature, the main diffraction peak shifts to lower diffraction angles. Using Bragg s law, we can determine the change of the lattice constant and, subsequently, the thermal expansion coefficient. The pronounced (111) orientations observed are in agreement with literature on Gd0.2Ce0.8O1.9-x wet-chemical and vacuum processed thin films 1 5. NATURE MATERIALS 1

2 Figure S1: In-situ heating XRD scans of pulsed laser deposited Gd 0.2 Ce 0.8 O 1.9-x film on Si 3 N 4 /Si substrates for the temperature range of 20 to 600 C shown at a heating rate of 5 C/min. S2: Evolvement of the lattice constant and thermal expansion coefficient during the insitu heating XRD for a Gd0.2Ce0.8O1.9-x substrate-supported thin film In Fig. S2, the lattice constant, a, and the thermal expansion coefficient, α, are plotted as obtained from the X-ray diffractograms with respect to the in-situ heating temperature (from Fig. S1). We report a lattice constant of around 5.6 Å for the (111) main diffraction line, which increases to 5.64 Å when heated from room temperature to 600 C at a constant rate of 5 C/min; this agrees well with a recent report placing the lattice constant at around 5.43 Å 1. While increasing the temperature we observed an increase in the lattice constant by 0.71% for the substrate-supported films due to the thermal expansion of the lattice, see Fig. S2a. The thermal expansion coefficient α at different temperatures is also reported in Fig. S2b. The thermal expansion coefficient, as determined from the (111) diffraction line at 600 C, is α 600 = C -1 and is close to literature values of, for example, Gd0.2Ce0.8O1.9-x pellets, which have a value of about C NATURE MATERIALS

3 SUPPLEMENTARY INFORMATION Figure S2: Lattice constant and thermal expansion coefficient of the substrate-supported Gd 0.2 Ce 0.8 O 1.9-x thin film with respect to the selected XRD diffraction lines and heating: (a) the lattice constant, a, and (b) the thermal expansion coefficient,α. S3: Oxygen ionic conductivity of the strained and free-standing Ce 0.8 Gd 0.2 O 1.9-x membrane: Conductivity measurements as a function of oxygen partial pressure and temperature We probed the dominant charge carrier type, ionic vs. electronic, for the strained Ce 0.8 Gd 0.2 O 1.9-x membranes via an oxygen partial pressure- and temperature-dependent electrical conductivity measurement carried out in a custom-made microprobe station setup, Fig. S3a, b. Through oxygen partial pressure and temperature-varied conductivity measurements, information on the prevailing dominant charge carrier, i.e. whether ionic conduction σ i or electronic conduction σ e, can be obtained. The total conductivity can be described by tot i e H H E (2 q)[ VO ] exp( ) e n exp( ) exp( ) T kt kt T kt 0i m 0e h 0 0 B B B (1) NATURE MATERIALS 3

4 where T denotes the absolute temperature, 2q the charge of an oxygen vacancy, [ V O ] their concentration, ν oi the pre-exponential ionic mobility factor containing geometrical factors and the jump attempt frequency, ΔH m the enthalpy of oxygen vacancy migration, and k B the Boltzmann constant. e 0 the electron charge, ν e the electronic mobility, n 0 the pre-exponential factor of electron concentration, ΔH the enthalpy of oxygen extraction, ν oe the pre-exponential electronic mobility factor, and E h the electron hopping energy. It is important to note that only the electronic conductivity is dependent on the oxygen partial pressure: The oxygen non-stoichiometry, i.e. such as the formation of additional oxygen vacancies and electrons during reduction of the gadolinia-doped ceria, is balanced by the gaseous phase of the measurement, see equation (2) where m represents the oxygen partial pressure exponent. As a consequence, potential variations of the total conductivity with respect to the varied oxygen partial pressure would identify a predominant electronic conduction for an isothermal. 1 m e T po (2) 2 In contrast, the dominant oxygen ionic carrier conductivity was reflected by an oxygen partial pressure independent conductivity at a given isothermal. Here, we show the result of the total conductivity measurement for the free-standing Ce 0.8 Gd 0.2 O 1.9-x buckled membrane with 0.46% compressive strain between electrode P2 and P4 at room temperature, Fig. S3b, as a function of the oxygen partial pressure and temperature in a Brouwer-type diagram, Fig. S3c, and an Arrhenius-type diagram, Fig. S3d. The total conductivity is independent of the oxygen partial pressure for the membranes tested in this study at temperatures of 300 to 500 C. 4 NATURE MATERIALS

5 SUPPLEMENTARY INFORMATION Based on this experimental evidence we unequivocally demonstrate that a dominant oxygen ionic conductivity prevails for the 0.46% compressively strained (between P2 and P4) and free-standing Ce 0.8 Gd 0.2 O 1.9-x film membranes. Additionally, our findings for the 0.46% compressively strained film membranes are in agreement with literature for Ce 0.8 Gd 0.2 O 1.9-x pellets and substrate-supported thin films, reporting a dominant oxygen ionic conductivity for the temperature range of C and higher oxygen partial pressures such as in air 7 9. Figure S3: Oxygen partial pressure and temperature-dependent measurements of the total conductivity of Ce 0.8 Gd 0.2 O 1.9-x thin film membranes with 0.46% compressively strain between P2 and P4. (a) Custom-made measurement setup. (b) Light microscopy image of a free-standing strained Ce 0.8 Gd 0.2 O 1.9-x thin film membrane with a 4-point electrode arrangement. (c) Log total electrical conductivity vs. log oxygen partial pressure Brouwer-type plot for isothermals. (d) Arrhenius-type diagram of the total conductivity with respect to the gas atmosphere. NATURE MATERIALS 5

6 S4: Description of the stress and strain states for substrate-supported thin films and free-standing Ce 0.8 Gd 0.2 O 1.9-x membranes: Thermodynamics and electro-chemomechanics Thermodynamics. Electro-chemo-mechanics describe the connection between oxygen ionic transport ( electro ), oxygen non-stoichiometry ( chemo ) and strained volumes ( mechanic ) of ionic conducting ceramics such as doped ceria 10. Here, the oxygen ionic conduction happens via the movement of ions hopping over oxygen vacancies. Additional external stress imposed on the crystal lattice can lead to volume and lattice position changes. Or, it can even induce the association and dissociation reactions between point defects in the material (e.g. dopant-oxygen vacancy associates 11,12 ). The formation of oxygen vacancies in the crystal lattice is governed by minimizing the Gibbs free energy, G. Accordingly, the change in free energy for spontaneity of defect formation has to be negative f vibr elastic config 0 G x H T S U T S (3) where ΔH f and U elastic are the enthalpy of formation of a defect and the mean elastic lattice energy, and ΔS vibr and ΔS config are the vibrational and configurational entropy due to the formation and arrangement of x defects in the atomic lattice In the case of elastic deformation acting on the ionic crystal lattice, an expression in terms of the stress tensor s ij, the strained cell volume V and the differential strain tensor ij is then given for the ionic conductor by 14 : du V s d (4) elastic ij ij ij For the cubic fcc unit cell of Ce 0.8 Gd 0.2 O 1.9-x 7, the cell volume can be expressed in terms of the lattice constant, i.e. V =a 3. For isotropic strain, this implies that measured volume changes can be directly converted to strain in one direction 16. The oxygen ionic conductivity, ionic, is 6 NATURE MATERIALS

7 SUPPLEMENTARY INFORMATION thermally activated by an oxygen ionic hopping mechanism over oxygen vacancy, [ V O ], and directly related to the strained lattice constant, a, of the material through (2 q) [ V ] (2 q) [ V ] (2 q) [ V ] ionic (5) kt kt kt O O 2 O 2 Dionic d a B B B 2q represents the charge of an oxygen vacancy, D ionic the ionic diffusion coefficient, d the jump distance, the ionic hopping frequency, k B the Boltzmann constant, and a geometry factor for the fcc cubic lattice of Ce 0.8 Gd 0.2 O x. The enthalpy of oxygen ionic migration, H mig, can be expressed thermodynamically by the Gibbs free energy (in accordance to equation (3)) and is directly correlated to the strained lattice volume and ionic diffusion using equations (4) and (5) to v H 0 mig ionic 2[ q VO ] vionic 2[ q VO ] exp( ) T kt B (6) where v ionic denotes the oxygen ionic mobility and v 0 is the pre-exponential ionic mobility factor. It can be concluded, based on a thermodynamic argument, that there is a direct impact between straining a cubic lattice and its point defect formation and the migration of oxygen ions, such as in the given example of Ce 0.8 Gd 0.2 O 1.9-x materials. We will use these basic arguments to discuss and define the stress and strain states for our model experiments on substrate-supported vs. free standing buckled Ce 0.8 Gd 0.2 O 1.9-x membranes in the following. Description of the Stress and Strain States for Substrate-Supported Thin Films and Free-Standing Ce 0.8 Gd 0.2 O 1.9-x Membranes. Thin films deposited on substrates are subject to biaxial residual stresses resulting from thermal expansion mismatch 18,19 and arising from film growth 20. The latter process is called atomic peening 18,19. This process describes atoms NATURE MATERIALS 7

8 incorporated in the thin film with higher density during deposition at high temperatures. Subsequently, in the course of cooling, the thin film relaxes to a lower density creating intrinsic stress 10,18,21. For example, it is reported that the biaxial residual stress in doped zirconia 18 and ceria 11,12,22 membranes can vary by 100s of MPa from compressive to tensile with respect to the aspect ratio of the film 18 or the deposition temperature 23 employed in stateof-the-art vacuum deposition techniques (i.e. pulsed laser deposition 24, and sputtering 21 ). It is this biaxial residual strain that defines what happens upon substrate removal. A deposited ionic conductor film will either crack due to tension or buckle under compression to a freestanding membrane 18,21,24. This is known in the field of micro-fuel cell membrane processing (see Refs. 18,21,23 27 for further details). We exemplify the overall net tensile and compressive strain (also called net strain ) acting on a Ce 0.8 Gd 0.2 O 1.9-x membrane leading to either cracking (tensile strain) or buckling (compressive strain), see Fig. S4a, b. Here, the difference in the initial intrinsic strain level was established through a change in the deposition temperature from 400 C (compressive strain) to 700 C (tensile strain) for the pulsed laser deposited Ce 0.8 Gd 0.2 O 1.9-x films. The following can be concluded on the thermodynamics of defect formation concerning the net tensile and compressive strain ( net strain ) of an ionic conducting thin film membrane: after substrate removal the Gibbs free energy is minimized to a new equilibrium state through buckling of the free-standing ceria-based membrane (see thermodynamic equations (3) and (4) for details). In this state the buckled thin film membrane is describable by a classic plate clamped onto a stiff substrate, e.g. the Karman plate model 28. Biaxial stress acts on the compressed thin film membrane for which we can define the strain tensor for in-plane (with i, j {x, y}) and outof-plane (z) components by () z () z (7) m b ( ij) ij 0 ij ij 8 NATURE MATERIALS

9 SUPPLEMENTARY INFORMATION where δ ij is the Kronecker delta, 0 is the average value of the residual strain, m is the in- ij b plane membrane strain and ( z) the out-of-plane bending and twisting modes of the ij membrane. We relate the out-of-plane deflection, ω, to the out-of-plane membrane strain, b ( z) of the thin film membrane from ij 2 b ij ( z) z (8) ij From the strain theorem of equation (8), the strain can be defined over the derivative of the out-of-plane deflection ω (function of buckling curve) of a compressed Ce 0.8 Gd 0.2 O 1.9-x membrane. We refer to the maximum strain change as the net compressive strain (1 st order strain) between the electrodes on the thin film membranes. The two cases for 1 st order strain are presented in Fig. S5a. In Fig. S5b, we schematically display the related out-of-plane amplitude, ω, for the free-standing film membrane. Additionally, we provide an example for the net compressive strain between the microelectrodes of a free-standing Ce 0.8 Gd 0.2 O 1.9-x membrane using a micrograph in Fig. S6. The elastic energy, U elastic (see equations (3) and (4)), of the in-plane Ce 0.8 Gd 0.2 O 1.9-x membrane components can be defined for the compressed film volumes using Hooke s law by 24 : Eh Eh U dxdy dxdy xy 2(1 ) 1 a/2 a/2 a/2 a/ elastic 2 ( xx yy 2 xx yy ) a/2 a/2 ( 0) a/2 a/2 z 0 z (9) where the terms a and h are side length and thickness, respectively, of the thin film, E is Young s modulus and ν is Poisson s ratio, see Fig. S5b. In this first part we describe and define the net compressive strain for the free-standing Ce 0.8 Gd 0.2 O 1.9-x membranes. This refers to the 1 st order maximum strains measured in between the locality of the microelectrodes. It is important to note that we also observed a 2 nd order waviness that superimposes this in-plane net compressed strain at smaller amplitudes for the Ce 0.8 Gd 0.2 O 1.9-x membranes, see Fig. S5c and Fig. S6. The overall net compressed NATURE MATERIALS 9

10 membrane reveals alterations of curvature as hilltops and valleys, see Fig. S6. Various techniques confirm the existence and variation of these 2 nd order local strain changes over various length scales (i.e. atomistic near order changes by Raman spectroscopy and macroscopic curvature changes by optical profilometry). Based on Stoney s equation 29, these local 2 nd order strain changes that superimpose the overall net compressively strained Ce 0.8 Gd 0.2 O 1.9-x membranes were segmented and computed using 2 hs Ek s 61 v h s GDC (10) where σ is the in-plane stress component in the film, k is curvature, E s is Young s modulus of the substrate, ν s is Poisson s ratio of the substrate, h s is thickness of the Si-substrate, and h GDC is the thin film thickness. Figure S4: Cracking or buckling of free standing Ce 0.8 Gd 0.2 O 1.9-x membranes caused by net tensile or compressive strain through different deposition temperatures. (a) The Ce 0.8 Gd 0.2 O 1.9-x membrane cracks after 10 NATURE MATERIALS

11 SUPPLEMENTARY INFORMATION PLD deposition at 700 C and free-etching of the initially substrate-supported film. (b) A buckled and compressively strained membrane occurs when the initial PLD deposition temperature is changed to 400 C. Figure S5: Strain states in free standing membranes. (a) Schematic figure of the isotropic tensile or compressive strain s influence on free standing Ce 0.8 Gd 0.2 O 1.9-x membranes, leading to either cracking or buckling, respectively. (b) Schematic description of a buckling membrane with microelectrodes. The net compressive strain is defined by the maximum out-of-plane deflection, ω max, of the compressed membrane. (c) Nomenclature of local strain based on the 2 nd order waviness, where one can observe hilltops and valleys. Figure S6: Example for the nomenclature of the 1 st (red) and 2 nd (orange) order strain based on a light microscopy image of a freestanding Ce 0.8 Gd 0.2 O 1.9-x membrane. NATURE MATERIALS 11

12 S5: Impedance spectroscopy of free standing Gd0.2Ce0.8O1.9-x electrolyte membranes The impedance spectra are measured for the Gd0.2Ce0.8O1.9-x electrolyte films and membranes. One membrane with the microelectrode design a is exemplified in the impedance response with respect to temperature, Fig. S7a. Here, we report RC-semicircles composed of a high frequency response, representing ionic transport in the membrane, grain and grain boundary, which are comprised of overlapping frequencies. The low frequency might be related to the electrode reaction and exchange with the gas phase: In Fig. S7b, an appliance of up to 1-3 V DC bias reveals that the spectra remains unchanged in the high frequency response at 500 C. Figure S7: Nyquist plots ofa Gd0.2Ce0.8O1.9-x electrolyte free-standing membrane measured with microelectrodes of design a. (a) Electrochemical impedance response with respect to temperature and (b) electrochemical impedance response with respect to applied DC bias. S6: Near order Raman spectroscopy of the Gd 0.2 Ce 0.8 O 1.9-x free standing membranes and substrate-supported thin films: Characterizing the oxygen anionic-cation near order Raman spectra of Gd 0.2 Ce 0.8 O 1.9-x free standing membranes and substrate-supported thin films are presented in the main text. For the self-supported membrane, Raman measurements with a spot size of ~850 nm were carried out at different localities of the free-standing membrane, namely the hilltops and valleys. As indicated in Fig. S8, these arrows indicate 12 NATURE MATERIALS

13 SUPPLEMENTARY INFORMATION predominant tensile and compressive local in-plane strain, respectively. The relationship between in-plane and out-of-plane strain is explained more detailed in the recent paper by Schweiger 17. The spectra for valleys reveal 4 peaks, all of them originating from the Gd 0.2 Ce 0.8 O 1.9-x : at ~250 cm -1, ~464 cm -1, ~557 cm -1, and ~600 cm -1. Literature analysis reveals that these can be ascribed to the following Raman features: second order transversal acoustic mode, 1 st order allowed F 2g mode, 2 nd order longitudinal acoustic and 2 nd order transversal optical mode 1,15,30,31. In the substrate-supported sample, another peak is visible at ~300 cm -1, which can be attributed to 2 nd order transversal acoustic phonon contributions of the crystalline silicon 32. The same assignment of peak modes is successful for the hilltops spectra, but with altered effective positions for the F 2g mode as described in the main text, Fig. 2h. Due to the penetration depth of the laser used for excitation we also measure a portion of the signal from the opposite side of the membrane, resulting in peak broadening. The asymmetry in the Raman signature is attributed to phonon dispersion and confinement effects due to the small grain size and strains in the materials 33. Figure S8: Raman measurements and peak position comparison. The arrows indicate the direction of the incident beam. Measurements were carried out on the hilltops where tensile strain prevails and valleys where compressive strain is predominant. NATURE MATERIALS 13

14 S7: Arrhenius diagrams of free standing membranes vs. substrate-supported thin films for all types of microelectrode designs In Fig. 3d of the main text, the activation energy determined from the ionic transport measurement is reported for the various electrode designs for the flat and substrate support films of Gd0.2Ce0.8O1.9-x and their compressively strained membrane counterparts. Here, in Fig. S9, the Arrhenius-type diagrams for those combinations of electrodes are supplemented in detail: Fig. S9a and S9b exhibit the plots for the free standing membranes and substratesupported thin films, respectively. Generally, one can observe that the ionic conductivity is higher for the substrate-supported thin films with lower activation energy, when compared to strained membranes under compression. Figure S9: Arrhenius diagrams for the measured conductivity in air for (a) free standing membranes and (b) substrate-supported thin films of Gd 0.2 Ce 0.8 O 1.9-x. Measurements were carried out for the different microelectrode designs a-c during heating (closed symbols) and cooling (open symbols). S8: Visualization of strain values in free standing membrane In Fig. 4 of the main text, the result of the strain distribution determined by the wafer curvature technique is viewed from the top of the membrane. Figure S10 shows the same 14 NATURE MATERIALS

15 SUPPLEMENTARY INFORMATION result viewed from a different angle to give an indication of the observed range of strain values. From this plot it can be deduced that the strains of the membrane locally range from -2% to 2% except for a number of outliers found at the edge of the membrane. Figure S10: Strain distribution in free-standing ionic conducting Ce 0.8 Gd 0.2 O 1.9-x electrolyte film determined by wafer curvature technique. Here, the membrane is viewed from the side with all scan lines included. The color code indicates the magnitude of measured strains. S9: The process of fabrication of free standing electrolyte membranes with microelectrodes Figure S11 illustrates the process of microfabrication for the free-standing membrane samples. The process is a combination of microfabrication and chemical wet etching and is detailed in the caption of Figure S11. NATURE MATERIALS 15

16 Figure S11: Manufacturing process flow processing of Gd 0.2 Ce 0.8 O 1.9-x free-standing membranes. (a) Cleaning of the Si substrate, (b) Si 3 N 4 layers are coated onto the two sides of the silicon wafer by Low Stress-CVD (LPCVD), (c) photolithography is employed to define the area to be etched via Reactive Ion Etching (RIE) for underneath the Si 3 N 4 layers to open space to the silicon, (d) KOH wet etching is used to remove the Si substrate so that a free-standing Si 3 N 4 membrane forms, (e) PLD deposition is used to deposit a Gd 0.2 Ce 0.8 O 1.9-x layer, (f) a 2 nd RIE step is used to etch the Si 3 N 4 underneath the Gd 0.2 Ce 0.8 O 1.9-x, forming now a free-standing membrane, (g) Pt films are deposited and structured by a shadow mask and e-beam evaporation and (h) the top microelectrodes are contacted via micropositioners in a custom-made microprobe station and attached to electrochemical test equipment as detailed in the methods section. 16 NATURE MATERIALS

17 SUPPLEMENTARY INFORMATION References 1. Rupp, J. L. M. et al. Scalable Oxygen-Ion Transport Kinetics in Metal-Oxide Films: Impact of Thermally Induced Lattice Compaction in Acceptor Doped Ceria Films. Adv. Funct. Mater. 24, (2014). 2. Mohan Kant, K., Esposito, V., & Pryds, N. Strain induced ionic conductivity enhancement in epitaxial Ce 0.9 Gd 0.1 O 2 δ thin films. Appl. Phys. Lett. 100, (2012). 3. Suzuki, T., Kosacki, I., & Anderson, H. Microstructure electrical conductivity relationships in nanocrystalline ceria thin films. Solid State Ionics 151, (2002). 4. Rodrigo, K. et al. The effects of thermal annealing on the structure and the electrical transport properties of ultrathin gadolinia-doped ceria films grown by pulsed laser deposition. Appl. Phys. A 104, (2011). 5. Rupp, J. L. M. Ionic diffusion as a matter of lattice-strain for electroceramic thin films. Solid State Ionics 207, 1 13 (2012). 6. Hayashi, H., Kanoh, M., Quan, C., & Inaba, H. Thermal expansion of Gd-doped ceria and reduced ceria. Solid State Ionics 132, (2000). 7. Mogensen, M., Sammes, N., & Tompsett, G. Physical, chemical and electrochemical properties of pure and doped ceria. Solid State Ionics 129, (2000). 8. Rupp, J. L. M., Infortuna, A. & Gauckler, L. J. Thermodynamic Stability of Gadolinia- Doped Ceria Thin Film Electrolytes for Micro-Solid Oxide Fuel Cells. J. Am. Ceram. Soc. 90, (2007). 9. Rodrigo, K. et al. Electrical characterization of gadolinia-doped ceria films grown by pulsed laser deposition. Appl. Phys. A 101, (2010). 10. Tuller, H. L. & Bishop, S. R. Point Defects in Oxides: Tailoring Materials Through Defect Engineering. Annu. Rev. Mater. Res. 41, (2011). 11. Lubomirsky, I. Practical applications of the chemical strain effect in ionic and mixed conductors. Monatshefte für Chemie - Chem. Mon. 140, (2009). 12. Lubomirsky, I. Stress adaptation in ceramic thin films. Phys. Chem. Chem. Phys. 9, (2007). 13. Atkins, P. & Paula, J. Physical Chemistry, Ninth Edition. Oxford Univ. Press (2010). 14. Maier, J. Physical Chemistry of Ionic Materials: Ions and Electrons in Solids. John Wiley Sons (2004). 15. Kossoy, A. et al. Influence of Point-Defect Reaction Kinetics on the Lattice Parameter of Ce 0.8 Gd 0.2 O 1.9. Adv. Funct. Mater. 19, (2009). NATURE MATERIALS 17

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19 SUPPLEMENTARY INFORMATION 31. Weber, W., Hass, K. & McBride, J. Raman study of CeO 2 : second-order scattering, lattice dynamics, and particle-size effects. Phys. Rev. B 48, (1993). 32. Wang, R. et al. Raman spectral study of silicon nanowires: High-order scattering and phonon confinement effects. Phys. Rev. B 61, (2000). 33. Dohčević-Mitrović, Z. D. et al. The size and strain effects on the Raman spectra of Ce 1 x Nd x O 2 δ (0 x 0.25) nanopowders. Solid State Commun. 137, (2006). NATURE MATERIALS 19