Fuel cells convert chemical energy directly into electrical energy with high efficiencies and environmental benefits, as compared with traditional heat engines1,2,3,4. Yttria-stabilized zirconia is perhaps the material with the most potential as an electrolyte in solid oxide fuel cells (SOFCs), owing to its stability and near-unity ionic transference number5. 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 membranes6. 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 SOFC 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.
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Financial support was provided by the Army Research Office (grants W911NF-14-1-0348 and W911NF-14-1-0669), the Air Force Office of Scientific Research (grant FA9550-12-1-0189), 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, Office of Science, Office of Basic Energy Sciences, under contract number DE-AC02-06CH11357. D.D.F. was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, Materials Sciences and Engineering Division.
The authors declare no competing financial interests.
Extended data figures and tables
Extended Data Figure 1 The electronic structure of SNO and H-SNO in the covalent limit and their electronic transport mechanisms.
a, The electronic structure of SNO in the covalent limit. Ligand holes are present on O 2p orbitals of the pristine SNO, while two electrons occupy the Ni eg manifold. The pristine SNO, however, is not strongly correlated because carriers transport through the O 2p ligand holes. b, The electronic structure of H-SNO. 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 SNO. c, The resistivity ρ of H-SNO compared with pristine SNO. The resistivity of H-SNO is more than eight orders of magnitude larger than that of pristine SNO at room temperature. d, e, Derivatives of resistivity (−dlnρ/dlnT) as a function of T plotted in log–log scale for H-SNO and SNO. The transport mechanism can be determined from the slope p of the −dlnρ/dlnT versus T curves. H-SNO shows the Efros–Shklovskii variable range hopping mechanism (p = 1/2), indicating polaron formation in the presence of a Coulomb gap31 (d). Pristine SNO shows crossover from activated conduction (p = 1) to Mott variable range hopping (p = 1/4) (e). The Coulomb repulsion is less strong in pristine SNO.
a, Temporal evolution of SNO conductivity when switching between different gas environments at various temperatures. b–d, Images of SNO and H-SNO on transparent substrate LAO. b, Pristine SNO shows dark, shining colour and the Pt bars are bright. c, After annealing in 5% H2/95% Ar at 300 °C for 1.5 h and cooling down to room temperature in the same gas environment, SNO 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 SNO indicates a diffusion profile of protons from the triple phase boundaries. The diffusion length LD is estimated to be ~300 μm.
Extended Data Figure 3 A schematic of the fabrication process of fuel cells with free-standing SNO membranes as the electrolyte.
a, Patterning a etch mask on the back side of the Si3N4/Si/Si3N4 chip with photoresist (PR). b, Removing exposed silicon nitride by reactive ion etching in CF4 and O2. c, Etching the Si from the back side with a KOH aqueous solution to make free-standing Si3N4 membrane. d, Depositing SNO thin films onto the Si3N4 membranes by radio-frequency magnetron sputtering and post-annealing the sample to form stoichiometric SNO. 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 SNO, 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.
a, An image of a 10 mm × 10 mm Si3N4/Si chip with nine SNO-electrolyte fuel cells (US dime coin shown for size). b, c, Optical micrographs of the free-standing buckled SNO 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 SOFC. d, A scanning electron microscope of the top porous Pt electrode. e, A schematic of the customized low-temperature micro-fabricated SOFC (μSOFC) testing station. Both pure H2 and 5%H2/95% Ar were used as fuel in the experiments.
a, Temporal evolution of the OCV of a Pt/H-SNO/Pt micro-fabricated SOFC with a 3% humidified 5% H2/95% Ar as fuel and laboratory air as oxidant, as the temperature ramps up. Initially SNO is electronically conductive so the OCV is close to zero. During the hydrogenation process, the OCV continues to increase after the temperature is stabilized and reaches near-ideal OCV, indicating that electronic conduction is almost completely suppressed in H-SNO by the Mott transition. The hydrogen fuel was always supplied at a constant flow rate both before t = 0 and during the experiments, and the initial low OCV is not due to the lack of fuel. b, The ionic transference number of H-SNO at various temperatures of interest to low-temperature SOFCs measured in Pt/SNO/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 OCV condition (infinitely large external resistance load) is modelled with an equivalent circuit containing a voltage source with an output voltage of Nernst potential EN, and two resistors Rion and Re, 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 Rpolarization element (drawn in red). VOC is the measured OCV. Note that there will be a small leakage current ileak 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 (Rpolarization) 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).
a, The dependence of micro-fabricated SOFC performance on the thickness of the SNO electrolyte at 500 °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 Ohmic resistance is varied while the electrode polarization resistance is kept more or less a constant. A clear increase in OCV with increasing thickness can be seen, which could be due to the decrease in the electrode polarization loss because of the larger electrolyte Ohmic 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 Ohmic resistance, but also higher OCV. b, Performance of Pt/SNO/Pd micro-fabricated SOFCs with a dense Pd anode with 3% humidified pure H2 as fuel and laboratory air as oxidant. It has been shown that hydrogen primarily creates protonic defects rather than oxygen vacancies in SNO (ref. 12). To verify that protons are the dominant mobile ion species in SNO and H-SNO, we fabricated an SOFC with the SNO electrolyte, a dense 100-nm-thick Pd anode, and a porous 100-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 SNO. During the fuel cell testing, 100 sccm pure H2 was flowed on the anode side, with the cathode exposed to air. The fuel cell with dense Pd has an OCV of 0.6 V and a peak power density of 24 mW cm−2 at 500 °C. The protonic conductivity of H-SNO can be extrapolated from impedance spectroscopy and OCV 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.
a, Cell voltage measured at 500 °C for a Pt/SNO/Pd fuel cell with wet 100% H2 as the fuel and stationary air as oxidant with current being 0 mA cm−2 (OCV condition) and 78 mA cm−2, respectively. The operation is stable for more than 20 h, implying that H-SNO 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 500 °C. b, X-ray diffraction pattern of SNO, and H-SNO (on LAO substrates) after being annealed under 1 bar of pure H2 at 500 °C for 48 h and 72 h. No new diffraction peaks are observed after annealing, which shows that H-SNO is quite stable in pure H2 for extended periods of time. θ is the incident angle of the X-ray.
a, Ex situ angle-dependent XANES spectra of hydrogenated SNO with a reference spectrum from pristine SNO. The critical angle θc of X-ray scattering for SNO at the X-ray energy near Ni K-edge is calculated to be 0.335°. When the X-ray incident angle is below the critical angle (0.25°), the XANES signal is surface sensitive with a penetration depth of ~10 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 0.25° 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.
a, An increase in the lattice constant can be caused by the larger crystal radius of Ni2+ and electron localization. When the formal valence state of Ni reduces, its ionic radius RNi increases, leading to the elongation of the Ni–O bond. In addition to the simple valence-state-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〉 = 0, 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 length32, that is, 〈Ni–O〉loc exceeds 〈Ni–O〉itin even for the same valence state. b, An optical image of a hydrogenated SNO sample. H-SNO 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 SNO and LAO peaks are indexed in pseudocubic notation. As the pristine SNO has a pseudocubic lattice constant close to that of the LAO, the SNO (002) appears almost as a shoulder of the LAO (002) peak. With decreasing distance between the X-ray spot and Pt electrodes, SNO (002) indeed shifts to smaller qz (no other peaks observed). Two peaks (peak 1 at qz = 3.18 Å−1 and peak 2 at qz = 2.98 Å−1) appear in the hydrogenated region and correspond to ~4% and ~10% 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 SNO (Extended Data Fig. 10c). d–f, Real-space mapping of the intensity of the Pt (111) peak at qz = 2.78 Å−1 (d), the H-SNO peak 1 at qz = 3.18 Å−1 (e) and the H-SNO peak 2 at qz = 2.98 Å−1 (f). A clear positive correlation between the Pt (111) and the qz = 2.98 Å−1 peaks can be seen, whereas the Pt (111) and qz = 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–O bond length can be also inferred from XANES spectra using Natoli’s rule33, which states that the energy separation between features B, D, and E will scale inversely with the square of the Ni–O distance, because they are derived from the first oxygen coordination cell27.
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 50 μm (horizontal in Extended Data Fig. 9b) × 500 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 qz = 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-SNO is still epitaxial on LAO 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 qz and qx 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 SNO directly underneath the porous Pt electrodes is on average closer to the triple phase boundaries (TPBs) than the SNO 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 ~ 100 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 LAO.
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Zhou, Y., Guan, X., Zhou, H. et al. Strongly correlated perovskite fuel cells. Nature 534, 231–234 (2016). https://doi.org/10.1038/nature17653
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