Rational Design of Perovskite-based Anode with Decent Activity for

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Rational Design of Perovskite-based Anode with Decent Activity for Hydrogen Electro-oxidation and Beneficial Effect of Sulfur for Promoting Power Generation in Solid Oxide Fuel Cells Yufei Song, Wei Wang, Jifa Qu, Yijun Zhong, Guangming Yang, Wei Zhou, and Zongping Shao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b11871 • Publication Date (Web): 01 Nov 2018 Downloaded from http://pubs.acs.org on November 2, 2018

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ACS Applied Materials & Interfaces

Rational Design of Perovskite-based Anode with Decent Activity for Hydrogen Electro-oxidation and Beneficial Effect of Sulfur for Promoting Power Generation in Solid Oxide Fuel Cells Yufei Song†, Wei Wang‡*, Jifa Qu†, Yijun Zhong‡, Guangming Yang†, Wei Zhou†, Zongping Shao†,‡* †Jiangsu

National Synergetic Innovation Center for Advanced Material, State Key Laboratory of

Materials-Oriented Chemical Engineering, College of Chemical Engineering, Nanjing Tech University, Nanjing, 210009, China ‡WA

School of Mines: Minerals, Energy and Chemical Engineering (WASM-MECE), Curtin

University, Perth, WA 6845, Australia Corresponding Authors * E-mail: [email protected] (W. W.). * E-mail: [email protected] (Z. S.).

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ABSTRACT: The poor sulfur tolerance of conventional nickel cermet anodes is particularly concerning for solid oxide fuel cell (SOFC) technology. Herein, we report an innovative anode composed of a samaria-doped ceria (SDC) scaffold and a perovskite La0.35Ca0.50TiO3-δ (LCT) thin film with a surface modified with strongly coupled and in situ formed Ni nanoparticles; the anode was prepared via an infiltration-calcination-reduction method. The rational design of such an anode transforms the detrimental effect of sulfur on the cell performance (poisoning) of state-of-the-art Ni cermet anodes into a beneficial effect promoting power generation from H2. A cell with a Ni+SDC cermet anode and a Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF) cathode showed an 18.3% reduction in the power output at 800 °C when the fuel gas was switched from pure H2 to H2-1000 ppm H2S, while a similar cell with this innovative anode showed an power output enhancement of 6.6%. Furthermore, the operational stability was significantly improved. The perovskite phase was found to account for the improved cell power output in the presence of the sulfur impurity. The introduction of the nickel nanoparticles further significantly enhanced the electrode activity, while the strong coupling effect of exsolved nickel nanoparticles with the perovskite thin film improved the sulfur tolerance of the nickel phase. As a result, the anode showed both high activity and stability while operating on H2 fuel with high concentration of H2S (1000 ppm). The promoting effect of sulfur on the power generation over the perovskite anode is also discussed. KEYWORDS: solid oxide fuel cells; anode; sulfur tolerance; beneficial effect; perovskite oxide 1. INTRODUCTION Solid oxide fuel cells (SOFCs) that allow the direct conversion of chemical energy stored in a chemical substance (fuel) into electric power via high-temperature electrochemical reactions over their electrodes are a relatively new technology characterized by high energy conversion efficiency, low emissions and excellent fuel flexibility.1-6 Low-temperature fuel cells require pure hydrogen (H2) as the fuel; thus, hydrocarbons should first be converted to H2 through steam reforming and the water gas shift reaction, which reduces the overall efficiency of the fuel cells. Due to the high operational temperatures, hydrocarbon fuels can be directly fed into the anode chamber of the SOFCs for power generation, and 2 Environment ACS Paragon Plus

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the power generation is realized via direct electrochemical oxidation or indirectly by the in situ conversion of hydrocarbon into H2 and CO over the anode (internal reforming). As a result, SOFCs have the highest overall energy efficiency among the various types of fuel cells when fuelled with hydrocarbons. In addition, compared with pure H2 fuel, the hydrocarbons are more attractive and practical due to their abundance and easy storage. This means that SOFCs may be ready for wide-spread application if their operational stability is satisfied since all the currently available infrastructure of energy delivery can be easily transferred to the SOFC technology.7-9 Over the past decade, much research has shown that nickel is a preferred electrocatalyst for fuel electro-oxidation in SOFCs due to its superior activity, conductivity and thermal compatibility.3 However, one challenge presented by the state-of-the-art SOFCs with Ni-based anodes for operation on hydrocarbon fuels derived from fossil fuels is the rapid performance degradation due to the sulfur poisoning of the anode. As we know, the hydrocarbons derived from fossil fuels always contain a certain level of sulfur impurity, even after deep refining. The sulfur adsorption easily occurs over the nickel surface to create a strong chemical bond with the formation of nickel sulfide, which blocks the adsorption of hydrogen for the subsequent electro-oxidation reaction.10,11 In extensive research studies, three different strategies have been widely exploited to mediate or solve the sulfur poisoning issue of the SOFC anode, i.e., the development of alternative anode materials, the optimization of the operation mode of the SOFCs, and the designing of special electrode architectures for minimizing the deposition of sulfur.10,12-16 For example, replacing the conventional electrolyte in the anode by the proton conductor BaZr0.1Ce0.7Y0.1Yb0.1O3-δ, introducing water vapour into the fuel and modifying the Ni-based anode with some oxygen or water storage materials could improve the sulfur tolerance of the anode.12-15 The origin of such improvements is that the oxidation rate of sulfur to gaseous SO2 is enhanced in the presence of such water/oxygen storage materials which allows for the re-exposure of the nickel surface for the electrochemical oxidation of H2 or other fuels for power generation. Another way to increase the sulfur tolerance of the nickel catalysts is to create a strong coupling effect between the nickel and the substrate,



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which may change the electronic structure of the nickel and reduce the bond formation energy between the nickel and sulfur.17 In addition to Ni-based anodes, considerable research activities have been conducted towards the development of novel materials with intrinsically higher sulfur tolerance.18,19 Among the various types of oxide materials, perovskite oxides are promising anode candidates due to their excellent chemical stability and superior carbon/sulfur tolerance.20-24 Various perovskite oxides such as Sr2Mg1-xMnxMoO6,

La0.7Sr0.3VO3- and La0.75Sr0.25Cr0.5Mn0.5O3- have been studied as the anode materials for SOFCs.21-23

Among the various perovskite oxides, titanate-based perovskite materials such as SrTiO3- and BaTiO3- are highly attractive due to their thermal and chemical stability in reducing atmospheres containing H2S.25,26 Nanostructure construction and morphology control are used to improve the sulfur tolerance of perovskite-based SOFC anodes.27 For example, the core-shell structured Li0.33La0.56TiO3 (LLTO) perovskite anode showed superior sulfur tolerance and operational stability compared to a nanostructured LLTO anode due to its excellent mechanical stability.27 However, all the abovementioned strategies may just result in a limited improvement in sulfur tolerance and, more importantly, may simultaneously impair other important parameter, such as the activity for fuel oxidation. For example, for the bulk-phase perovskite oxide anode, although the sulfur tolerance was significantly improved, the power output of the fuel cells with those anodes remained poor due to their relatively poor activity for fuel electro-oxidation reactions at intermediate temperatures.21,22,25-30 This suggests that new anode materials/configurations are still urgently needed to realize simultaneous high power output and superior sulfur tolerance. Herein, we report a novel perovskite-based anode with a superior sulfur tolerance, high electrocatalytic activity and robust stability designed based on the combination of a core-shell structure and an in situ exsolution with a limited phase transition (stable substrate). By controlling the firing temperature for the infiltrated anode, a core-shell structured La0.33Ca0.47Ti0.94Ni0.06O3-δ (LCTN3)/Sm0.2Ce0.8O1.9 (SDC) was fabricated. After a subsequent treatment in a reducing atmosphere at 800 4 Environment ACS Paragon Plus

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°C, a strongly coupled hierarchical architecture composed of nickel nanoparticles and core-shell structured La0.35Ca0.50TiO3-δ (LCT)/SDC substrate was formed. Instead of poisoning the anode, such perovskite shows an improved power output from H2 containing sulfur impurity (H2S). The reason for the superior performance of the new anode and the surprisingly enhanced performance in the presence of sulfur for power generation over the perovskite oxide are discussed. 2. EXPERIMENTAL SECTION 2.1. Materials The SDC electrolyte, BSCF and Sm0.5Sr0.5CoO3-δ (SSC) cathodes were prepared by a sol-gel method.27 The as-synthesized precursors were then sintered at 700, 950 and 1000 °C for 5 h, respectively. The LCT and La0.33Ca0.47Ti1-xNixO3-δ (LCTN, x=0.02, 0.04, 0.06 and 0.08, which are denoted as LCTN-1, LCTN-2, LCTN-3 and LCTN-4, respectively) were prepared via a solution combustion method and the precursors were calcined at 850 °C in air for 2 h. The proper amount of deionized water and stoichiometric amounts of La(NO3)3•6H2O, Ca(NO3)2•4H2O, Ti(OC3H7)4, Ni(NO3)2•6H2O were dissolved together to prepare Ni, LCT and LCTN precursors for infiltration. The concentration of metal ion was 0.4 mol L-1. Citric acid was added to promote the dissolution of tetrabutyltitanate. Glycine was also added as a complexing agent in the infiltration process. The Ba(Zr0.4Ce0.4Y0.2)0.8Ni0.2O3-δ (BZCYN) precursor for infiltration was prepared by the same method.17 2.2. Cell fabrication The single cells were prepared as follows. First, the dense SDC pellets were prepared by drypressing and calcining at 1400 °C for 5 h in air. The diameter of the as-synthesized SDC pellet was approximately 11.9 mm. The appropriate SDC powder, isopropyl alcohol and 10 wt.% soluble starch (a pore former) were mixed by ball milling to form a suspension that was sprayed onto one surface of the as-synthesized SDC pellet with an effective surface area of 0.45 cm2. The spray-painted electrode was sintered at 1250 °C for 5 h in air to yield the porous SDC scaffold with a thickness of ~30 μm. The BSCF and SSC cathode solutions were then sprayed onto the other surface of the electrolyte in the same 5 Environment ACS Paragon Plus


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way with a thickness of ~20 μm and calcined at 1000 °C in air for 2 h. The as-prepared LCTN precursor solutions were infiltrated into the SDC scaffold and detailed information can be found in the literature.17 The mass fraction of the infiltrated LCT or LCTN in the anode was 30 wt.%. Afterwards, the composite anode was sintered at 850 °C for 2 h. For the purpose of comparison, similar steps were also utilized to prepare Ni, LCT and BZCYN-infiltrated SDC anodes. To create a reliable comparison, the mass fraction of Ni (~1 wt.%) in the Ni-containing anodes was equal to the LCTN infiltrated SDC anode. The pure LCT, LCTN, and NiO+SDC (3:2, by weight) suspensions, which were prepared by ball milling, were sprayed onto one surface of the as-synthesized SDC pellets with an effective surface area of 0.45 cm2 and calcined at 850, 850, and 1250 °C for 5 h, respectively. Silver paste was coated on the electrode surface as a current collector. 2.3. Characterization The Keithley 2420 source metre was used to get the I-V and I-P curves of the fuel cells with a fourprobe configuration. The anode chamber was fed by H2 or H2-1000 ppm H2S fuels with a flow rate of 80 mL min-1 [STP] and the cathode was exposed to ambient air. The impedance of the fuel cell was surveyed by an electrochemical impedance spectroscopy (EIS) measurement using a Solartron 1260 frequency response analyser and a Solartron 1287 potentiostat under the open circuit voltage (OCV) condition. In the stability test, the SSC cathode replaced the BSCF cathode due to the better stability of the former cathode. To prevent the volatilization of silver during the long-term stability test, the Ag paste was replaced by an Ag mesh-like current collector.17 The phase structures of the as-synthesized powders were determined using X-ray diffraction (XRD, D8 Advance Bruker) with a Cu-K source. The microstructures of the LCTN-infiltrated electrodes were characterized by a field emission scanning electron microscope (FE-SEM, JEOL-S4800) and a highresolution transmission electron microscope (HR-TEM, JEOL JEM-2100). Bright-field scanning transmission electron microscopy (STEM) images were obtained using an FEI Tecnai G2 T20 electron microscope. The corresponding energy-dispersive X-ray (EDX) mappings were obtained by a TEMEDX apparatus (FEI Tecnai G2 F30 STWIN). The TEC data was collected using Netzsch DIL 6 Environment ACS Paragon Plus

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402C/3/G dilatometer. X-ray photoelectron spectroscopy (XPS) measurements were performed on a Thermo ESCALAB 250 using monochromatic Al Kα radiation (1486.6 eV). The BET surface areas of the LCTN samples before and after H2 reduction were analysed by nitrogen adsorption-desorption isotherms using a Quanta chrome AutoSorb-iQ3 instrument at liquid nitrogen temperature. The H2 temperature-programmed reduction (H2-TPR) was monitored by an in situ thermal conductivity detector (TCD) with a BELCAT-A apparatus. 17 3. RESULTS AND DISCUSSION First, the maximum Ni doping amount at the B-site of LCT perovskite oxide was studied. Figure 1 shows the XRD patterns of the Ni-free LCT and various LCTN powders calcined at 850 °C for 2 h. The NiO phase was only detected in the as-synthesized LCTN-4 sample, suggesting that the maximum Ni doping amount at the B-site of LCT perovskite could not exceed 8 mol.% for the formation of a pure perovskite phase. Then, we tested the power output and operational stability of the SDC electrolytesupported single cell with the pure LCT or LCTN anodes operating with H2 and H2-1000 ppm H2S fuel to choose the best anode candidate. As shown in Figure 2a and Figure S1a, the peak power densities (PPDs) of the fuel cells with the LCT, LCTN-1, LCTN-2 and LCTN-3 anodes operating with H2 fuel were 6, 14, 28 and 39 mW cm-2 at 800 °C, respectively. Furthermore, the LCTN-3 anode exhibited the best operational stability with H2 fuel among the three LCTN anodes (Figure S1b). The cell performances in terms of PPDs and polarization resistances of the single cells with LCT or LCTN anodes operating with H2-1000 ppm H2S fuel was also tested, as shown in Figure S2 and Figure S3. The PPDs of the fuel cell with LCT, LCTN-1, LCTN-2 and LCTN-3 anodes operating on H2-1000 ppm H2S fuel were 25, 57, 91 and 123 mW cm-2, respectively, at 800 oC (Figure 2a). The enhancement factor, defined as the ratio of the PPDs of the pure LCT/LCTN anode in H2-1000 ppm H2S to their PPDs in H2 fuel, is shown in the inset of Figure 2a. The LCT anode showed the highest enhancement factor for PPD when the fuel was switched from H2 to H2-1000 ppm H2S. However, the absolute value of the PPD was still too low. The nickel doping improved the performance of the LCT anode with a slight reduction in the enhancement factor for the PPD. Based on the above experimental results, the LCT and 7 Environment ACS Paragon Plus


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LCTN anodes showed significantly improved power outputs in H2 fuel with H2S, suggesting that the H2S promoted the anodic reaction of the LCT-based perovskite oxides. Furthermore, it was found that the pure LCTN-3 anode displayed the best operational stability and sulfur tolerance in H2-1000 ppm H2S at 800 °C (Figure 2b), which indicates it be the most promising SOFC anode candidate of these four materials. However, the electrocatalytic activity and stability of the pure LCTN-3 anode was still too poor for the potential commercialization. Infiltration and morphology control are typically used to improve the electrocatalytic activity and operational stability.16,27 Therefore, we developed a LCTN-3/SDC composite anode through infiltration and calcination. For the composite electrode, the phase structure and morphology are two important characteristics that should be addressed. First, the phase structures of the LCTN-3 powder before and after treatment in H2 were investigated (Figure 3). As shown in Figure 3a, the LCTN-3 sample was well indexed to an orthorhombic perovskite structure with lattice parameters of a=b=5.433Å and c=7.706 Å. After H2 treatment, a weak diffraction peak at a 2-theta of 44.5° appeared, which was assigned to the (110) diffraction peak of Ni, indicating the formation of a metallic nickel phase. In addition, the main orthorhombic perovskite phase was maintained, suggesting the LCTN-3 perovskite has high structural stability, and could provide a stable substrate for a nickel exsolution under reducing atmosphere. The lattice parameters for the main perovskite were found to be a=b=5.457Å and c=7.744Å for the sample after the H2 reduction, which are larger than those of the fresh LCT before the reduction (a=b=5.425 Å, c=7.694 Å), which is attributed to the Ni exsolution and the partial Ti4+ reduction (Figure 3b). The ionic radius of Ti3+ is 0.670 Å, which is larger than that of Ti4+ (0.605 Å). This suggests that the following reaction is likely take place during the H2 reduction process: La0.33Ca0.47Ti0.94Ni0.06O3-δ+0.06H2=La0.35Ca0.50TiO3-δ+0.06Ni+0.06H2O.


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Figure 1. XRD patterns of LCT, various LCTN samples with different Ni doping amounts.

Figure 2. a) The PPDs of LCT, LCTN-1, LCTN-2, LCTN-3 anodes operating on H2 and H2-1000 ppm H2S at 800 °C. Inset is the comparison of the performance enhancement factors of various anodes. b) The stability tests of the fuel cells with LCTN-1, LCTN-2, LCTN-3 anodes operating on H2-1000 ppm H2S under a current density of 50 mA cm-2 at 800 °C.



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Figure 3. a) XRD patterns of LCT, LCTN-3 and reduced LCTN-3 samples. b) Ti 2p XPS spectra of the LCTN-3 sample before and after treatment in H2. To further support the exsolution of nickel from the LCTN-3 perovskite lattice and examine the morphological shape of the resulting metallic nickel, SEM, HR-TEM and STEM combined with EDX spectroscopy were conducted. The surface of the unreduced LCTN-3 was smooth (Figure S4), while many nanoparticles with diameters of approximately 15 nm appeared after the treatment in H2 at 800 °C (Figure 4a). From the TEM image shown in Figure 4b, some nanoparticles with diameters of approximately 15-20 nm were found on the surface of the main grains. As shown in Figure 4c, the nanoparticles and the main grains showed diffraction fringes with the lattice spacing of 0.180 and 0.221 nm, which were assigned to the (002) and (022) diffraction planes of the metallic Ni and LCT phases, respectively. The Ni exsolution was also confirmed by the EDX results as shown in Figure 4d. This result suggests that the main grains are composed of a perovskite phase with lattice parameters that are the same as LCT while the nanoparticles are metallic nickel in nature. From the elemental mapping images shown in Figure 4e, the reduced LCTN-3 particles are composed of La, Ca, Ti and Ni, and Ni shows a very different distribution compared to the other elements, suggesting the exsolution of Ni nanoparticles. Figure 4f and g show the particle-substrate interface for the exsolved Ni particle and the Ni nanoparticle is socketed into the perovskite substrate, which is consistent with the literature.31,32 The nickel lattice growing from the LCTN-3 perovskite lattice might naturally facilitate the interdiffusion between the metal lattice and the oxide lattice, which has been shown shown to significantly increase 10 Environment ACS Paragon Plus

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the adhesion between the metal and oxide phases even when occurring over thin 1-2 unit cell interfaces. Thus, a strong interaction between Ni nanoparticles and LCT substrate is realized.

Figure 4. a) SEM, b) TEM and c) HR-TEM images of the reduced LCTN-3 powder. d) EDX results of Ni nanoparticle. e) STEM-EDX results of reduced LCTN-3 powder. f) TEM image of reduced LCTN-3 powder and g) the particle-substrate interface for exsolved Ni nanoparticle. The phase structure and morphology of the composite anode fabricated by infiltration were also characterized. Figure 5a shows the XRD patterns of the as-prepared SDC, LCTN-3, and LCTN-3infiltrated SDC anode before and after H2 reduction. For the unreduced LCTN-3-infiltrated SDC sample, the XRD patterns matches well with that of the physical mixture of the SDC and LCTN-3 phases, suggesting that phase reaction between the SDC scaffold and LCTN-3 electrocatalyst are negligible. After the H2 treatment, the diffraction peaks of the composite anode were largely unchanged, suggesting that the main perovskite structure survived after the treatment in the reducing atmosphere. The SEM images of the porous SDC scaffold and the LCTN-3-infiltrated SDC electrode before and after H2 reduction are shown in Figure 5b-d. The SDC scaffold is highly porous and well-distributed 11 Environment ACS Paragon Plus


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particles (200-400 nm) as shown in Figure 5b. After the introduction of the LCTN-3 phase by infiltration and sintering at 850 °C, the inner surface of the SDC scaffold was found to be decorated with a LCTN-3 thin film (Figure 5c). To further verify the formation of the core-shell structured LCTN-3/SDC, STEM and EDX were performed and the results are shown in Figure S5. The elements of Sm and Ce were mainly distributed throughout the centre of the imaged particle and La, Ca and Ti were mainly distributed around the edges of this particle (Table S1), suggesting the formation of a coreshell structure. After the reduction in H2, some smaller Ni nanoparticles approximately 15 nm in size formed on the surface of the infiltrated perovskite thin film, as shown in Figure 5d.

Figure 5. a) XRD patterns of SDC, LCTN-3, LCTN-3-infiltrated SDC and LCTN-3-infiltrated SDC after H2 treatment. SEM images of the b) SDC scaffold, c) thin-film LCTN-3-infiltrated SDC anode and d) thin-film LCTN-3-infiltrated SDC anode after H2 treatment. Figure 6 shows the TEM, HR-TEM and linear EDX scanning images of one SDC particle modified by thin-film LCTN-3 and Ni nanoparticles after the infiltration/reduction processes. The surface of the SDC grain was decorated with a porous layer composed of Ni nanoparticles and LCT perovskite oxide (Figure 6a). The TEM image and linear EDX analysis clearly reveal the core-shell structure of the thinfilm LCTN-3-infiltrated SDC anode (Figure 6b). As shown in Figure 6c, the exsolved Ni nanoparticles 12 Environment ACS Paragon Plus

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are also socketed into the thin-film LCT perovskite substrate, which is consistent with the TEM result of the reduced LCTN-3 powder (Figure 4g). The surface area of the electrodes in the SOFCs plays a critical role in the electrode activity since most of the SOFC electrodes are heavily sintered. Even a small enhancement in the surface area could benefit the electrode reaction due to the accelerated gas diffusion inside the electrode and the increased triple phase boundary length. The nickel nanoparticles increased the specific surface area from 14.04 to 16.01 m2 g-1 after the reduction (Figure S6), which is consistent with the literature.33

Figure 6. a) TEM images of reduced thin-film LCTN-3-infiltrated SDC anode and linear EDX scanning of Ni nanoparticle, b) linear EDX scanning of the core-shell particle and c) HR-TEM images of reduced thin-film LCTN-3-infiltrated SDC anode. Electrical conductivity plays a significant role in the SOFC electrode performance. The conductivity of the reduced thin-film LCTN-3-infiltrated SDC in 10 vol.% H2-Ar is shown in Figure 7a. This anode display a conductivity of 0.12-0.43 S cm-1 in the temperature range between 600 and 800 °C, which is comparable with the values reported for perovskite oxide-based anodes.27,34-37 A good thermo-mechanical compatibility between the cell components is critical for achieving a high 13 Environment ACS Paragon Plus


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operational stability for the SOFCs. One significant advantage of the infiltrated electrodes in SOFCs is that thermal expansion behaviour is similar to that of the electrolyte material, ensuring superior durability for SOFCs. Figure 7b shows the thermal expansion curves of the SDC scaffold and the thinfilm LCTN-3-infiltrated SDC anode after the reduction. The TEC of the SDC scaffold was calculated to be 12.6×10-6 K-1 at 300-800 °C, which is close to the value reported in the literature.17 For the reduced thin-film LCTN-3-infiltrated SDC, the TEC for the same temperature range was 12.4×10-6 K-1. The similarity in the TEC values suggests that there is a good compatibility between the reduced thin-film LCTN-3-infiltrated SDC electrode and the SDC electrolyte, which provides for a good thermomechanical compatibility and operational stability.

Figure 7. a) The conductivity of the reduced thin-film LCTN-3-infiltrated SDC anode in 10 vol.% H2Ar at the temperature range between 550-800 °C. b) TEC curves of the SDC scaffold and reduced thinfilm LCTN-3-infiltrated SDC anode in Ar. The performance tests of SOFCs with H2 and H2-H2S as fuels were conducted to evaluate the electrochemical activity of the reduced thin-film LCTN-3-infiltrated SDC anode. An electrolytesupported single cell consisting of a thin-film LCTN-3-infiltrated SDC anode, SDC electrolyte and BSCF cathode was fabricated, as shown in Figure S7, and their thicknesses were approximately 30, 400 and 20 μm, respectively. As shown in Figure 8a, the single cell with the reduced thin-film LCTN-3infiltrated SDC anode exhibited PPDs of 443, 348, 257 and 164 mW cm-2 at 800, 750, 700 and 650 °C, 14 Environment ACS Paragon Plus

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respectively, when fuelled by H2. Considering that the single cell was fabricated with an SDC electrolyte with a thickness of 400 μm, the PPDs achieved in this study are highly attractive. As shown in Figure 8b, when fuelled by H2-1000 ppm H2S, the single cell delivered PPDs of 474, 360, 278 and 173 mW cm-2 at 800, 750, 700 and 650 °C, respectively, which are higher than those of the H2S-free H2. The EIS results shown in Figure 8c and d further confirm the promoting effect of the H2S in H2 fuel for the reduced thin-film LCTN-3-infiltrated SDC anode. The enhanced cell performance that occurs in the presence of H2S suggest that H2S promoted the H2 oxidation and improved the performance, which is consistent with the pure LCTN-3 anode. The way that the H2S adsorb on the anode plays an important role in enhancing the performance with H2S-containing fuels. The chemisorption reaction between sulfur and oxygen to form [H2SO]+ occurs very easily even at room temperature.8,38 The [H2SO]+ transports proton to the oxygen through H-S-O and releases of 787.6 kJ mol-1 of energy.38 In contrast, the dissociation of H-H bond to produce proton is an endothermic reaction requiring -435.8 kJ mol-1 of energy.39 Therefore, transferring the proton to the oxygen through H-S-O is much easier than breaking the H-H bond. Furthermore, adsorbed sulfur regenerates easily in H2 to produce H2S. In addition, H2S is also a hydrogen carrier in the fuel, which promotes the transportation of H2 from the gas to the surface of the anode and, therefore, oxidation of H2 was promoted.



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Figure 8. The I-V and I-P curves of the fuel cells with reduced thin-film LCTN-3-infiltrated SDC anode operating on a) H2 and b) H2-1000 ppm H2S at 650-800 °C. The EIS spectra of the fuel cells with reduced thin-film LCTN-3-infiltrated SDC anode operating on c) H2 and d) H2-1000 ppm H2S at 650800 °C. The EIS spectra of the fuel cells with e) pure LCTN-3 and f) thin-film LCTN-3-infiltrated SDC anode operating on H2 and H2-1000 ppm H2S at 800 °C. However, the performance enhancement of the reduced thin-film LCTN-3-infiltrated SDC anode was much smaller than that of the pure LCTN-3 anode, which could be explained by the EIS analysis. For the pure LCTN-3 anode, the performance was very low due to the poor activity and oxygen ion conductivity of the anode. Additionally, the electrode resistance (Rp) played a predominant role in the total resistance (Rtotal), as shown in Figure 8e. The introduction of H2S in the fuel could promote the H2 oxidation to decrease Rp. The reduction in Rp would result in the significant enhancement of 16 Environment ACS Paragon Plus

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performance. For the reduced thin-film LCTN-3-infiltrated SDC anode, the performance was higher than that of the pure LCTN-3 anode due to the higher activity and oxygen ion conductivity of the infiltrated anode. However, the ohmic resistance (Ro) played a predominant role in the Rtotal, as shown in Figure 8f. Although the introduction of H2S in the fuel promoted H2 oxidation to reduce the Rp, the enhanced performance of the infiltrated anode was still smaller than that of the pure LCTN-3 anode. For comparison, the reduced thin-film LCTN-3-infiltrated SDC, Ni-infiltrated SDC, LCT-infiltrated SDC, Ni+SDC and BZCYN-infiltrated SDC were also investigated as anodes for SOFCs operating with H2 and H2-1000 ppm H2S fuels at 800 °C as shown in Figure 9 and Figures S8-S10. For H2 fuel, the PPDs of the cells with the reduced thin-film LCTN-3-infiltrated SDC, Ni-infiltrated SDC, LCTinfiltrated SDC, Ni+SDC and BZCYN-infiltrated SDC anodes were 443, 456, 41, 493 and 538 mW cm-2 (Figure 9a), respectively, while the corresponding PPDs with H2-1000 ppm H2S fuel were 474, 372, 90, 403 and 446 mW cm-2 (Figure 9b), respectively. The comparison of the performances of the different anodes is shown in Figure 9c. Compared with the PPDs of the cells fuelled by H2, the PPDs of the fuel cells fuelled by H2-1000 ppm H2S and the Ni-infiltrated SDC, Ni+SDC and BZCYN-infiltrated SDC anodes were reduced by 18.4, 18.3 and 17.1%, respectively. On the other hand, the reduced thin-film LCTN-3-infiltrated SDC anode displayed a performance enhancement of 6.6% (Figure 9d), suggesting that the reduced thin-film LCTN-3-infiltrated SDC anode has excellent sulfur tolerance. Compared with the pure Ni-infiltrated SDC anode, the performance enhancement of the reduced thin-film LCTN-3infiltrated SDC anode should be attributed to the thin-film LCT, which has been shown to improve the cell performance after switching the H2 fuel to H2-1000 ppm H2S fuel. The EIS results shown in Figure 9e and f further confirm the superior electrocatalytic activity and sulfur tolerance of the reduced thinfilm LCTN-3-infiltrated SDC compared with the other three anodes.



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Figure 9. The I-V and I-P curves of the fuel cells with different anodes operating on a) H2 and b) H21000 ppm H2S at 800 °C. c) The PPDs of different anodes operating on H2 and H2-1000 ppm H2S at 800 °C. d) The performance enhancement of different anodes after switching the H2 to H2-1000 ppm H2S fuel at 800 °C. The EIS spectra of the fuel cells with different anodes operating on e) H2 and f) H2-1000 ppm H2S at 800 °C. The EIS spectra of the fuel cells with different Ni-containing anodes operating on H2 and H2-1000 ppm H2S at 800 °C is enlarged as shown in the inset figure. The LCTN-3-SDC, NiSDC and BZCYN-SDC represented for the infiltrated anodes and the Ni+SDC anode represented for the anode prepared by physical mixing. In addition to the power density, the operational stability also plays a vital role in establishing the commercialization of the SOFCs technology. Figure S11 shows the stability test of the fuel cells with the reduced thin-film LCTN-3-infiltrated SDC and Ni+SDC anodes operated with H2 fuel at a polarization current density of 100 mA cm-2 at 800 °C. As seen, the voltage of the fuel cell with the thin-film LCTN-3-infiltrated SDC anode remained stable at 0.72 V for 50 h, which is comparable to the 18 Environment ACS Paragon Plus

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Ni+SDC anode. Figure 10 shows the stability test of the fuel cells with the Ni-infiltrated SDC, Ni+SDC and reduced thin-film LCTN-3-infiltrated SDC anodes operated with H2-1000 ppm H2S at a polarization current density of 100 mA cm-2 at 800 °C. As seen, the Ni-infiltrated SDC and Ni+SDC anodes are not stable and the voltage decreased rapidly after ~1.2 and ~40 hours’ operation with H2-1000 ppm H2S fuel, respectively. The decay rates of the fuel cells with the reduced thin-film LCTN-3-infiltrated SDC, Ni-infiltrated SDC, and Ni+SDC anodes were 0.0018, 0.1384 and 0.0072 V h-1, respectively. In the long-term operational stability test, H2S may possibly be adsorbed on the surface of Ni to form Ni2S3, which easily dissolves at 800 °C since its melting point is only 787 °C.11 Therefore, the Ni-containing anodes exhibit different levels of performance degradation. However, the decay rates of Ni-infiltrated SDC and Ni+SDC anodes are higher than that of the thin-film LCTN-3-infiltrated SDC anode. This may be attributed to the weaken interaction between Ni and SDC substrate in the Ni-infiltrated SDC and Ni+SDC anodes, which lead to a stronger bond between the Ni nanoparticles and sulfur, and resulted the severe performance degradation. However, the exsolved Ni nanoparticles have a strong interaction with the thin-film LCT substrate, which can change the electronic structure of the Ni nanoparticles. Thus, the performance degradation was suppressed when the thin-film LCTN-3-infiltrated SDC was used as the anode as shown in Figure 10. In addition, the microstructure of the reduced thin-film LCTN-3-infiltrated SDC anode was practically unchanged after the stability test and some of the Ni nanoparticles were still found on the LCT thin film (Figure S12). These results highlight the superior stability and excellent sulfur tolerance of the anode based upon the core-shell structured substrate.



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Figure 10. The stability test of the fuel cells with Ni-infiltrated SDC, Ni+SDC and reduced thin-film LCTN-3-infiltrated SDC anodes operating on H2-1000 ppm H2S under a current density of 100 mA cm-2 at 800 °C. Strengthening the interaction between the Ni particles and the substrates could promote the improvement of sulfur/coking resistances of the anodes for SOFCs.31 We hypothesize that the reduced thin-film LCTN-3 perovskite oxide with exsolved Ni nanoparticles should have a much stronger interaction between the Ni nanoparticles and the substrate than the Ni-infiltrated SDC, Ni+SDC and BZCYN-infiltrated SDC anodes. To test this hypothesis, the as-prepared thin-film LCTN-3-infiltrated SDC, Ni-infiltrated SDC, Ni+SDC and BZCYN-infiltrated SDC anodes were subjected to H2-TPR analysis (Figure 11). The H2-TPR curve of NiO is also provided for comparison. The thin-film LCTN3-infiltrated SDC anode shows a much higher interaction than the pure NiO, Ni-infiltrated SDC, Ni+SDC and BZCYN-infiltrated SDC anodes, suggesting a stronger interaction between Ni nanoparticles and the substrate in the thin-film LCTN-3-infiltrated SDC anode. The differences in the interactions between Ni and the substrates of the various anodes can partially explain the different sulfur tolerances of the various anodes as shown in Figure 11.


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Figure 11. H2-TPR profiles of the Ni-infiltrated SDC, NiO+SDC, thin-film LCTN-3-infiltrated SDC and BZCYN-infiltrated SDC anodes and NiO. In addition to H2, many hydrocarbons, such as methane, ethanol, propane, etc., are often used as fuels for SOFCs. Carbon dioxide (CO2) is often produced in the hydrocarbon reforming reaction or introduced as a reaction additive in hydrocarbon reforming.40,41 Therefore, we further investigated the CO2 tolerance of the reduced thin-film LCTN-3-infiltrated SDC anode by testing the performance of the single cell operated with 20% CO2-H2 and 20% CO2-H2-1000 ppm H2S fuels at 650-800 °C. The single cell exhibited the PPDs of 436, 339, 240 and 150 mW cm-2 at 800, 750, 700 and 650 °C, respectively, when fuelled by 20% CO2-H2 (Figure 12 and Figure S13a). To evaluate the performance of the reduced thin-film LCTN-3- infiltrated SDC anode in CO2- and H2S-containing fuels, we further tested 20% CO2H2-1000 ppm H2S as the fuel for the SOFCs. The single cell produced PPDs of 460, 355, 263 and 160 mW cm-2 at 800, 750, 700 and 650 °C, respectively (Figure 12 and Figure S13b), when fuelled by 20% CO2-H2-1000 ppm H2S fuel, which are higher than those of the H2S-free 20% CO2-H2. This result is consistent with the thin-film LCTN-3-infiltrated SDC anode operated on CO2-free H2 or H2-1000 ppm H2S fuel. The results suggest that the thin-film LCTN-3-infiltrated SDC anode could be a promising candidate for SOFCs operated with hydrocarbon fuels.



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Figure 12. The PPDs of the fuel cells with reduced thin-film LCTN-3-infiltrated SDC anode operating on H2, 20% CO2-H2 and 20% CO2-H2-1000 ppm H2S at 650-800 °C. 4. CONCLUSIONS In conclusion, we find a surprising promotion effect of sulfur in the H2 fuel for the LCT-based perovskite oxide anode and develop a new strategy to rationally design the LCTN-3/SDC composite anode with a high activity for hydrogen electro-oxidation and excellent sulfur tolerance through infiltration and calcination. First, the promotion effect of sulfur in the H2 fuel for the LCT-based anode is found and the LCTN-3 anode exhibits the highest activity and stability operating with H2 and H2-1000 ppm H2S. However, the electrocatalytic performance of the LCTN-3 anode is still too poor for potential commercialization. Therefore, we try to develop a composite LCTN-3/SDC anode through the combination of infiltration and in situ exsolution. After calcining the LCTN-3-infiltrated SDC anode at 850 °C and treating it in H2 at 800 °C, a composite anode composed of an SDC scaffold and a La0.35Ca0.50TiO3-δ thin film with its surface modified with strongly coupled and in situ formed Ni nanoparticles is fabricated. This hierarchical composite anode displays chemical and thermal compatibility between the cell components as well as excellent electrochemical activity in H2 and H21000 ppm H2S fuels. As a result, this rationally designed anode displays higher PPDs than the Niinfiltrated SDC, BZCYN-infiltrated SDC and Ni+SDC anodes under the same conditions. In addition, 22 Environment ACS Paragon Plus

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the single cell with this composite anode deliver a superior operational stability for 65 h at 800 °C in H21000 ppm H2S fuel, which is greater than that of fuel cells with Ni-infiltrated SDC and Ni+SDC anodes under the same conditions. This composite anode also shows an excellent CO2 tolerance, which provides more possibilities for the application of this material in hydrocarbon-fuelled SOFCs. Our research offers a new strategy for developing perovskite-based electrocatalysts in the SOFCs to promote the practical application of the SOFC technology. Furthermore, this design strategy can be easily transferred to other electrocatalysis-based applications such as water splitting and metal-air batteries. Supporting Information The following information are available free of charge. Experimental section, additional figures, and tables as described in the text (PDF) Corresponding Authors * E-mail: [email protected] (W. W.). * E-mail: [email protected] (Z. S.). Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS This work was financially supported by National Natural Science Foundation of China of No. 21576135 and 21706129, Jiangsu Natural Science Foundation for Distinguished Young Scholars of No. BK20170043, and the Youth Fund in Jiangsu Province of No. BK20150945. The authors also acknowledge the financial support of the Australian Research Council via the Discovery Projects program (DP150104365 and DP160104835).

References 23 Environment ACS Paragon Plus


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Figure 1. XRD patterns of LCT, various LCTN samples with different Ni doping amounts. 222x180mm (300 x 300 DPI)



Figure 2. a) The PPDs of LCT, LCTN-1, LCTN-2, LCTN-3 anodes operating on H2 and H2-1000 ppm H2S at 800 oC. Inset is the comparison of the performance enhancement factors of various anodes. b) The stability tests of the fuel cells with LCTN-1, LCTN-2, LCTN-3 anodes operating on H2-1000 ppm H2S under a current density of 50 mA cm-2 at 800 oC. 278x117mm (96 x 96 DPI)


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Figure 3. a) XRD patterns of LCT, LCTN-3 and reduced LCTN-3 samples. b) Ti 2p XPS spectra of the LCTN-3 sample before and after treatment in H2. 501x221mm (96 x 96 DPI)



Figure 4. a) SEM, b) TEM and c) HR-TEM images of the reduced LCTN-3 powder. d) EDX results of Ni nanoparticle. e) STEM-EDX results of reduced LCTN-3 powder. f) TEM image of reduced LCTN-3 powder and g) the particle-substrate interface for exsolved Ni nanoparticle. 483x326mm (96 x 96 DPI)


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Figure 5. a) XRD patterns of SDC, LCTN-3, LCTN-3-infiltrated SDC and LCTN-3-infiltrated SDC after H2 treatment. SEM images of the b) SDC scaffold, c) thin-film LCTN-3-infiltrated SDC anode and d) thin-film LCTN-3-infiltrated SDC anode after H2 treatment. 284x203mm (96 x 96 DPI)



Figure 6. a) TEM images of reduced thin-film LCTN-3-infiltrated SDC anode and linear EDX scanning of Ni nanoparticle, b) linear EDX scanning of the core-shell particle and c) HR-TEM images of reduced thin-film LCTN-3-infiltrated SDC anode. 289x201mm (96 x 96 DPI)


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Figure 7. a) The conductivity of the reduced thin-film LCTN-3-infiltrated SDC anode in 10 vol.% H2-Ar at the temperature range between 550-800 oC. b) TEC curves of the SDC scaffold and reduced thin-film LCTN-3infiltrated SDC anode in Ar. 317x127mm (96 x 96 DPI)



Figure 8. The I-V and I-P curves of the fuel cells with reduced thin-film LCTN-3-infiltrated SDC anode operating on a) H2 and b) H2-1000 ppm H2S at 650-800 oC. The EIS spectra of the fuel cells with reduced thin-film LCTN-3-infiltrated SDC anode operating on c) H2 and d) H2-1000 ppm H2S at 650-800 oC. The EIS spectra of the fuel cells with e) pure LCTN-3 and f) thin-film LCTN-3-infiltrated SDC anode operating on H2 and H2-1000 ppm H2S at 800 oC. 348x357mm (96 x 96 DPI)


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Figure 9. The I-V and I-P curves of the fuel cells with different anodes operating on a) H2 and b) H2-1000 ppm H2S at 800 oC. c) The PPDs of different anodes operating on H2 and H2-1000 ppm H2S at 800 oC. d) The performance enhancement of different anodes after switching the H2 to H2-1000 ppm H2S fuel at 800 oC. The EIS spectra of the fuel cells with different anodes operating on e) H2 and f) H2-1000 ppm H2S at 800 oC. The EIS spectra of the fuel cells with different Ni-containing anodes operating on H2 and H2-1000 ppm H2S at 800 oC is enlarged as shown in the inset figure. The LCTN-3-SDC, Ni-SDC and BZCYN-SDC represented for the infiltrated anodes and the Ni+SDC anode represented for the anode prepared by physical mixing. 354x356mm (96 x 96 DPI)



Figure 10. The stability test of the fuel cells with Ni-infiltrated SDC, Ni+SDC and reduced thin-film LCTN-3infiltrated SDC anodes operating on H2-1000 ppm H2S under a current density of 100 mA cm-2 at 800 oC. 187x145mm (300 x 300 DPI)


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Figure 11. H2-TPR profiles of the Ni-infiltrated SDC, NiO+SDC, thin-film LCTN-3-infiltrated SDC and BZCYNinfiltrated SDC anodes and NiO. 173x135mm (300 x 300 DPI)



Figure 12. The PPDs of the fuel cells with reduced thin-film LCTN-3-infiltrated SDC anode operating on H2, 20% CO2-H2 and 20% CO2-1000 ppm H2S-H2 at 650-800 oC. 599x464mm (300 x 300 DPI)


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