Protonic ceramic fuel cells, like their higher-temperature solid-oxide fuel cell counterparts, can directly use both hydrogen and hydrocarbon fuels to produce electricity at potentially more than 50 per cent efficiency1,2. Most previous direct-hydrocarbon fuel cell research has focused on solid-oxide fuel cells based on oxygen-ion-conducting electrolytes, but carbon deposition (coking) and sulfur poisoning typically occur when such fuel cells are directly operated on hydrocarbon- and/or sulfur-containing fuels, resulting in severe performance degradation over time3,4,5,6. Despite studies suggesting good performance and anti-coking resistance in hydrocarbon-fuelled protonic ceramic fuel cells2,7,8, there have been no systematic studies of long-term durability. Here we present results from long-term testing of protonic ceramic fuel cells using a total of 11 different fuels (hydrogen, methane, domestic natural gas (with and without hydrogen sulfide), propane, n-butane, i-butane, iso-octane, methanol, ethanol and ammonia) at temperatures between 500 and 600 degrees Celsius. Several cells have been tested for over 6,000 hours, and we demonstrate excellent performance and exceptional durability (less than 1.5 per cent degradation per 1,000 hours in most cases) across all fuels without any modifications in the cell composition or architecture. Large fluctuations in temperature are tolerated, and coking is not observed even after thousands of hours of continuous operation. Finally, sulfur, a notorious poison for both low-temperature and high-temperature fuel cells, does not seem to affect the performance of protonic ceramic fuel cells when supplied at levels consistent with commercial fuels. The fuel flexibility and long-term durability demonstrated by the protonic ceramic fuel cell devices highlight the promise of this technology and its potential for commercial application.
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This work was supported by Advanced Research Projects Agency–Energy (ARPA-E) for funding under the REBELS program (award DE-AR0000493). Additional support was provided by the Army Research Office under grant no. W911NF-17-1-0051 and the Colorado Office of Economic Development and International Trade (COEDIT) under their Advanced Industries Proof-of-Concept grant program. The views and conclusions contained in this document are those of the authors and should not be interpreted as representing the official policies, either expressed or implied, of ARPA-E, the Department of Energy, the Army Research Office or the US Government. This research used resources of the Advanced Light Source, which is a DOE Office of Science User Facility under contract no. DE-AC02-05CH11231. The US Government is authorized to reproduce and distribute reprints for Government purposes notwithstanding any copyright notation herein.
Nature thanks J. Dailly, T. Norby, Z. Shao and the other anonymous reviewer(s) for their contribution to the peer review of this work.
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Extended data figures and tables
a, j–V and j–P curves of PCFC under hydrogen (flow rate of hydrogen: 15 ml min−1, cell 1). b, j–V and j–P curves of direct-methanol PCFC at 600–350 °C. Flow rate of the mixture of methanol and water (molar ratio of water/ethanol is 1, cell 3) is 0.024 ml min−1. c, j–V and j–P curves of PCFC under 10 vol% iso-C4H 10 + 20 vol% O2 + 70 vol% H2O (flow rate of iso-C4H10: 0.9 ml min−1, cell 4). d, j–V and j–P curves of PCFC under 10 vol% n-C4H10 + 20 vol% O2 + 70 vol% H2O (flow rate of n-C4H10: 0.9 ml min−1, cell 5). e, j–V and j–P curves of PCFC under 20 vol% C3H8 + 20 vol% O2 + 60 vol% H2O (flow rate of C3H8: 1.6 ml min−1, cell 4). f, j–V and j–P curves of PCFC under 33.3 vol% natural gas (simulated natural gas with 19.5 p.p.m. H2S, flow rate of natural gas: 5 ml min−1, cell 7) + 66.7 vol% H2O. g, j–V and j–P curves of PCFC under 33.3 vol% natural gas (simulated natural gas without H2S, flow rate of natural gas: 5 ml min−1, cell 7) + 66.7 vol% H2O. h, j–V and j–P curves of PCFC under 33.3 vol% methane (flow rate of methane: 5 ml min−1, cell 7) + 66.7 vol% H2O. i, j–V and j–P curves of PCFC under 28.5 vol% methane (flow rate of methane: 5 ml min−1, cell 8) + 71.5 vol% H2O. j, j–V and j–P curves of direct-ethanol PCFC at 600 °C and 550 °C. Flow rate of the mixture of ethanol and water (molar ratio of water/ethanol is 7, cell 10) is 0.005 ml min−1.
a, Direct-propane PCFC. b, Direct-iso-butane PCFC.
Extended Data Fig. 3 SEM images of PCFC (Cell 4) after running for about 6,000 h on hydrocarbons at 500 °C.
a, Anode (low magnification). b, Anode (high magnification). c, The sandwich structure of this cell after running for about 6,000 h. d, The interface between electrolyte and cathode. e, High magnification of cathode after 6,000 h operation. f, Raman spectrum of the PCFC anode after 6,000 h of operation on hydrocarbon fuel at 500 °C. A carbonate peak is visible at 1,060 cm−1 but the graphitic carbon (G band) and disordered carbon (D band), which would be present at 1,580 cm−1 and 1,350 cm−1, are not apparent, indicating the absence of graphitic carbon and disordered carbon deposits in the anode even after long-term operation on hydrocarbon fuel. Long-term stabilities of direct hydrocarbons PCFCs were tested at 500 °C. At this temperature, the structure type of carbon species should be polymeric, amorphous films or filaments (Cβ), vermicular filaments, fibres and/or whiskers (Cγ), or graphitic platelets or films (Cc)35. Typically, these carbon species are visible (by SEM) in SOFC anodes after running on hydrocarbon fuels, but the high-magnification SEM images of our PCFC anode (Extended Data Fig. 3a, b) show no visible evidence of such carbon structures. To further substantiate this conclusion, post-mortem Raman analyses show that there are no disordered and graphitic carbon species found in the PCFC anode even after long-term operation on hydrocarbon fuels (Extended Data Fig. 3f).
Extended Data Fig. 4 TEM image and EDS mapping of PCFC (Cell 3) after 8,000 h operation on methanol.
a, TEM image of lamella from BZY20 PCFC electrolyte/cathode interface prepared by FIB. b–h, HAADF image of the electrolyte/cathode interface, and corresponding EDS maps of Ba, Zr, Y, Co, Fe and O. There is no chemical reaction between the electrolyte and electrolyte which indicates the excellent chemical compatibility of cathode with electrolyte. EDS mapping of cathode shows there is no obvious elemental segregation.
Extended Data Fig. 5 Coking resistance of Ni/BZY and Ni/YSZ anodes investigated by in situ HT-Raman spectroscopy.
a, Raman spectra of Ni/YSZ anode exposed to humidified methane (S:C = 2; 500 °C) at various times up to 200 min. b, Raman spectra of Ni/BZY anode exposed to humidified methane (S:C = 2; 500 °C) at various times up to 360 min. c, Raman spectra of Ni/BZY anode exposed to humidified methane (S:C = 1; 500 °C) at various times up to 72 h (4,320 min). d, Raman spectra of Ni/YSZ anode exposed to humidified methane (S:C ≈ 1; 550 °C) at various times up to 40 min. e, Raman spectra of Ni/BZY anode exposed to humidified methane (S:C ≈ 1; 550 °C) at various times up to 100 min.
Extended Data Fig. 6 SEM image of the anode of a PCFC after running for 1,400 and 6,000 h on hydrocarbons at 500 °C.
a, b, Operation for 6,000 h on iso-butane and propane (cell 4). c, Operation for 1,400 h on methane. d, Operation for 6,000 h on iso-butane and propane (cell 4). The Tamman temperature of nickel is 581 °C. The target operating temperature of our PCFCs is lower than 600 °C, and most of the long-term stability testing conducted in this study was at 500 °C (that is, well below the Tamman temperature), at which Ni sintering/coarsening is avoided. Indeed, as shown in Extended Data Fig. 6c, d, the distribution and size of Ni nanoparticles on the BZY anode support remain essentially constant between 1,400 and 6,000 h of operation.
a, Exposure to reducing conditions (for example, H2 or hydrocarbon fuel) at typical PCFC operating temperatures triggers the in situ exsolution of Ni nanoparticles onto the surface of the BZY phase in the cermet anode. This exsolution process is driven by the decreased solubility of Ni in the BZY phase under anode operating conditions (500–600 °C, highly reducing) compared with the much higher starting solubility of Ni in the BZY phase under anode synthesis conditions (1,450 °C, highly oxidizing). In addition, fuel cell operation also drives the decomposition of a residual BaY2NiO5 phase which is a by-product of the solid-state reactive-sintering method (SSRS) used to synthesize the anode. As illustrated here, the SSRS process enables the anode to be rapidly and inexpensively fabricated in a single step starting from a homogeneous mixture of BaCO3, ZrO2, Y2O3, NiO and pore former, which is fired under oxidizing conditions at 1,450 °C for 18 h. During firing, a complex phase-formation and sintering process, involving the transient formation and decomposition of BaY2NiO5, leads to the creation of a porous two-phase NiO + Ni-doped BZY composite anode with a small amount of dispersed residual BaY2NiO5. During initial fuel cell operation, the NiO phase is reduced to Ni metal, while the Ni-doped BZY is reduced to BZY36, concomitant with the exsolution of Ni nanoparticles which we hypothesize enhance the performance and durability of the fuel cell. b, Amplified main peaks of XRD patterns of BZY20 phase after sintering in air with different NiO amount in precursors. The amount of Ni diffusing into the BZY is determined by the defect reaction equilibrium. The amplified main XRD peaks of BZY20 phase shift to the higher angles with increasing amounts of Ni in the anode. This is consistent with a decreasing lattice constant with increasing Ni substitution on the B-site of the BZY lattice owing to the smaller size of Ni2+ compared with Zr4+ or Y3+. c, XRD patterns of BZY20 anode after sintering in air and reduction in hydrogen at 600 °C for 150 h. d, Amplified main peaks of XRD patterns of BZY20 phase after sintering in air and reduction in hydrogen at 600 °C for 150 h. The peak shifts to smaller angles after reduction which is consistent with the exsolution of Ni from BZY lattice, leading to a concomitant increase in the lattice constant of the BZY phase. e, BaY2NiO5 phase formation temperature profile investigated by in situ high-temperature XRD. The colour indicates the intensity of the main peak of BaY2NiO5.
a, b, BZY20 phase of the anode before reduction. The surface of the BZY20 phase is very clean, and there are no nanoparticles along the grain boundaries or on the triple junction points. c, d, BZY20 phase of the anode after reduction under hydrogen at 600 °C for 50 h. Nickel nanoparticles begin to form on the triple junction points and grain boundaries. e, f, BZY20 phase of the anode after reduction under hydrogen at 600 °C for 100 h. More Ni nanoparticles begin to form along grain surfaces and boundaries. The size of exsolved Ni nanoparticles is less than 100 nm.
Extended Data Fig. 9 Evolution of the AP-XPS elemental survey spectrum of Ni/BaZr0.9Y0.1O3-δ (BZY) and Ni/Zr0.92Y0.08O2−δ (8YSZ) patterned microelectrodes.
a, Ni/BZY-based microelectrode, b, Ni/YSZ-based microelectrode.
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Duan, C., Kee, R.J., Zhu, H. et al. Highly durable, coking and sulfur tolerant, fuel-flexible protonic ceramic fuel cells. Nature 557, 217–222 (2018). https://doi.org/10.1038/s41586-018-0082-6
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