Highly Active and Stable Cobalt-Free Hafnium-doped SrFe0.9Hf0.1O3

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Article Cite This: ACS Appl. Energy Mater. 2018, 1, 2134−2142

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Highly Active and Stable Cobalt-Free Hafnium-doped SrFe0.9Hf0.1O3−δ Perovskite Cathode for Solid Oxide Fuel Cells Zhenbao Zhang,† Dengjie Chen,*,† Jian Wang,‡ Shaozao Tan,† Xiang Yu,† and Zongping Shao§ †

ACS Appl. Energy Mater. 2018.1:2134-2142. Downloaded from pubs.acs.org by UNIV OF EDINBURGH on 01/15/19. For personal use only.

Guangdong Engineering & Technology Research Centre of Graphene-like Materials and Products, Department of Chemistry, College of Chemistry and Materials Science, Jinan University, Guangzhou 510632, China ‡ Department of Mechanical and Aerospace Engineering, The Hong Kong University of Science and Technology, Hong Kong 999077, China § Jiangsu National Synergetic Innovation Center for Advanced Materials (SICAM), State Key Laboratory of Materials-Oriented Chemical Engineering, College of Chemical Engineering, Nanjing Tech University, Nanjing 210009, China S Supporting Information *

ABSTRACT: Sluggish oxygen reduction reaction (ORR) kinetics and chemical instability of cathode materials hinder the practical application of solid oxide fuel cells (SOFCs). Here we report a Co-free Hf-doped SrFe0.9Hf0.1O3−δ (SFHf) perovskite oxide as a potential cathode focusing on enhancing the ORR activity and chemical stability. We find that SFHf exhibits a high ORR activity, stable cubic crystal structure, and improved chemical stability toward CO2 poisoning compared to undoped SrFeO3−δ. The SFHf cathode has a polarization area-specific resistance as low as 0.193 Ω cm2 at 600 °C in a SFHf|Sm0.2Ce0.8O1.9 (SDC)|SFHf symmetrical cell and has a maximum power density as high as 1.94 W cm−2 at 700 °C in an anode-supported fuel cell (Ni+(ZrO2)0.92(Y2O3)0.08 (YSZ)| YSZ|SDC|SFHf). The ORR activity maintains stable for a period of 120 h in air and in CO2-containing atmosphere. The attractive ORR activity is attributed to the moderate concentration of oxygen vacancy and electrical conductivity, as well as the fast oxygen kinetics at the operation temperature. The improved chemical stability is related to the doping of the redox-inactive Hf cation in the Fe site of SrFeO3−δ by decreasing oxygen vacancy concentration and increasing metal−oxygen bond energy. This work proposes an effective strategy in the design of highly active and stable cathodes for SOFCs. KEYWORDS: Co-free, SrFeO3‑δ, structure stability, redox-inactive cation, Perovskite cathode

1. INTRODUCTION Fuel cells are recognized to be efficient and environmentally friendly energy conversion devices to mediate the world’s limited carbon resources and to reduce emissions.1−4 Among them, solid oxide fuel cells (SOFCs) can directly convert chemical energy stored in a wide variety of fuels into electricity without combustion. To make SOFCs commercialized, the high operating temperature (>800 °C) must be reduced. However, the decrease in operating temperature typically leads to the sluggish oxygen reduction reaction (ORR) kinetics at the cathode.1,2 Therefore, intensive efforts have been devoted to exploring novel compositions or architectures with high electrocatalytic activity toward ORR in order to realize commercialization.1,2,5−9 Perovskite-type oxides that possess mixed ionic and electronic conductivity have been considered to be the most promising cathode materials for the nextgeneration SOFCs. Perovskite-type mixed conducting electrodes are capable of significantly enlarging electrocatalytic active sites for ORR to whole electrode including inner and outer surfaces as thick as several micrometers, while traditional © 2018 American Chemical Society

electronic-conducting electrodes (e.g., La1−xSrxMnO3) with negligible ionic conductivity limit active sites for ORR to the restricted triple phase boundary line, where electrolyte, cathode, and oxygen gas meet.10 Considerable electrocatalytic activity toward ORR in the low- and intermediate-temperature range of 500−750 °C has been demonstrated primarily by Cocontaining perovskite-type oxides such as Ba1−xSrxCoyFe1−yO3−δ (BSCF),11,12 La1−xSrxCoO3−δ (LSC),13 Sm1−xSrxCoO3−δ,14 PrBaCo2O5+δ,15,16 and SrCo0.8Nb0.1Ta0.1O3−δ.8 For example, Zhou et al. modified the oxygen vacancy, oxygen ionic mobility, and electron transfer on a SrCo0.8Nb0.1Ta0.1O3−δ cathode, which exhibited excellent electrochemical performance such as a polarization area-specific resistance (ASR) as low as ∼0.16 Ω cm2 and a peak power density of ∼1.2 W cm−2 even at a relatively low temperature of 500 °C.8 Received: February 9, 2018 Accepted: April 11, 2018 Published: April 11, 2018 2134

DOI: 10.1021/acsaem.8b00198 ACS Appl. Energy Mater. 2018, 1, 2134−2142

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ACS Applied Energy Materials

oxygen permeability and the CO2 tolerance was taken seriously in our previous reports when developing perovskite-type materials for practical membrane applications. To the best of our knowledge, only a few studies have simultaneously investigated the activity and chemical stability in developing cathode materials for SOFCs.25,26 Herein, we report a Co-free Hf-doped perovskite oxide with a nominal composition of SrFe0.9Hf0.1O3−δ (SFHf) as cathode for SOFCs. Though HfO2 itself only shows an acidity scale similar to that of Fe2O3,27 it can change the acidity of other cations to a high acidity scale thanks to the high valence state of Hf4+. The successfully synthesized SFHf exhibits outstanding ORR activity and chemical stability (i.e., CO2 tolerance). A low ASR of 0.193 Ω cm2 is achieved at 600 °C when SFHf is assembled as electrodes in a symmetrical cell configuration. A peak power density as high as ∼1.94 W cm−2 at 700 °C is achieved in an anode-supported fuel cell. The attractive ORR activity is attributed to the moderate concentration of oxygen vacancy and electrical conductivity, and fast oxygen kinetics at the operation temperature. The improved chemical stability is due to the doping of the redox-inactive Hf cation in the Fe site of SF.

Unfortunately, these state-of-the-art Co-containing cathode materials suffer from activity degradation over time under operation conditions.17 Rupp et al. reported that ORR on the LSC cathode was either deactivated by intentionally depositing Sr on the surface of the electrode, or activated by modifying the surface with Co, suggesting that Sr element is detrimental while Co is beneficial for producing electrocatalytic active sites.18 Their results successfully supported the degradation mechanism, which was usually ascribed to the element segregation (e.g., Sr) on the surface, where insulating phases such as SrO blocked the active sites for surface electron transfer and oxygen.13 The element segregation may arise from positive and negative charge interactions and cation size mismatch of perovskites.19 To mediate this instability issue of Co-containing perovskite-type oxides such as LSC, it was proposed that coating La0.8Sr0.2MnO3 or ZrO2 thin films with several nanometers on the surface enhanced the ORR stability and activity,5 but the method applied for formation of thin films is time-consuming and complex for practical applications. Alternatively, Yildiz et al. reported that the oxygen surface exchange kinetics and stability of LSC could be simply enhanced upon decreasing the concentration of surface oxygen vacancy via depositing redox-inactive cations, e.g., Nb5+, Ti4+, Zr4+, Hf4+, and Al3+ on the LSC surface.13 Among the investigated compositions, Hf4+-modified LSC positioned the peak regarding the oxygen surface exchange rate.13 This suggests that both enhanced oxygen surface exchange kinetics and improved chemical stability could be realized by properly optimizing the surface oxygen vacancy concentration. The other critical issues of these state-of-the-art Cocontaining cathode materials often include the high cost of Co compared to that of Fe, thermal expansion mismatch with traditional electrolytes, poor crystal structure stability, high reactivity with ZrO2-based electrolyte, and susceptibility to CO2.1 Although the facile reduction of Co cations is beneficial to the creation of oxygen vacancy which improves the ORR activity, the coexistence of redox-active Co cations (e.g., Co2+, Co3+, and Co4+) should be mainly responsible for these critical issues of cathode materials. The CO2 tolerance is of particular importance for cathode working under ambient air or in singlechamber SOFCs since CO2 poisoning could destroy the crystal structure of the cathode materials and degrade the oxygen exchange kinetics.18 For example, the polarization ASRs of the SrSc0.175Nb0.025Co0.8O3−δ and BSCF cathodes obtained from symmetric cells increased by ∼20 times and ∼37 times after replacing 10 vol % N2 with CO2 at 600 °C.20 To enhance the CO2 tolerance of traditional cathodes, Zhou et al. developed a core−shell structure consisting of La2NiO4+δ dense thin films and cones on BSCF.21 However, the applied method is also time-consuming and the desired cathode architecture is hard to be precisely prepared. According to literature, the susceptibility to CO2 poisoning of perovskite-type oxides was closely related to the thermodynamics of carbonates, average metal−oxygen bond energy (ABE), Lewis acidity, and oxygen vacancy concentration.20,22,23 Recently a doping strategy has been widely employed to improve the CO2 tolerance. The picked dopants should possess characteristics of (1) high acidity, (2) high valence state, and (3) high ABE.24 Ti4+, Zr4+, Nb5+, Sb5+, Ta5+, W6+, and Mo6+ could probably help mediate the CO2 poisoning effects of perovskites.22−24 Very recently, we have found that Co-free high-valence and redox-inactive transition metal cations doped SrFeO3−δ (SF) membranes showed effectively improved CO2 tolerance.22,23 The balance of the

2. EXPERIMENTAL SECTION 2.1. Materials Preparation. The SrFe1−xHfxO3−δ (x = 0 and 0.1) oxides were synthesized using a high-energy ball-milling solid-state reaction method. SrCO3, Fe2O3, and HfO2 (Analytical grade) were used as raw materials. For a typical synthesis, stoichiometric amounts of raw materials were ball-milled (Fritsch, Pulverisette 6) with ethanol as the milling medium at a rotation speed of 400 rpm for 2 h. After drying, the primary powders were prefired at 1100 °C in air for 10 h, followed by calcining at 1250 °C for 10 h in air to obtain a desired perovskite oxide. Sm0.2Ce0.8O1.9 (SDC) powders were prepared by the sol−gel method applying ethylene diamine tetraacetic acid (EDTA) and citric acid (CA) as complexing agents and calcined at 800 °C for 5 h, and a detailed process can be found in our previous work.16 2.2. Symmetric Cell and Fuel Cell Fabrication. A symmetrical cell with a configuration of SFHf|SDC|SFHf was prepared for electrochemical impedance spectroscopy (EIS) measurement. First a dense SDC electrolyte membrane was fabricated using dry-pressing method, followed by sintering at 1400 °C for 5 h. Then the SFHf cathode powder was dispersed in a premixed solution of ethylene glycol, isopropyl alcohol, and glycerol to prepare the cathode slurry via high-energy ball-milling at a rotation rate of 400 rpm for 0.5 h. The obtained SFHf slurry was wet-sprayed symmetrically onto both surfaces of the SDC electrolyte membrane and then fired at 900 °C for 2 h in air. An anode-supported fuel cell with a configuration of Ni +(ZrO2)0.92(Y2O3)0.08 (YSZ)|YSZ|SDC|SFHf was fabricated for the electrochemical performance measurement. The anode-supported fuel cell was prepared via tape casting, spray deposition, and hightemperature sintering.7,28 The obtained SFHf slurry was wet-sprayed onto the central surface of the electrolyte side and calcined at 900 °C for 2 h in air. An effective surface area of the SFHf cathode is 0.45 cm2. 2.3. Basic Characterizations. The phase structure at room and elevated temperatures was examined by powder X-ray diffraction (XRD; Rigaku Smartlab 3 kW) in the 2θ range of 10−90° with a step size of 0.02°. The long-term phase structure stability was also measured by XRD. The detailed structural information was analyzed by performing Rietveld refinement on XRD patterns using GSAS software to obtain lattice parameters. The phase structure and elemental distribution of SFHf was characterized by high-resolution transmission electron microscopy (HR-TEM; JEOL JEM-2000) equipped with an energy-dispersive X-ray spectroscopy (EDS) detector. The chemical compatibility between SFHf and SDC was examined to investigate the phase reaction. The oxygen temperatureprogrammed desorption (O2-TPD; Hiden QIC-20) experiment and 2135

DOI: 10.1021/acsaem.8b00198 ACS Appl. Energy Mater. 2018, 1, 2134−2142

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ACS Applied Energy Materials thermogravimetric (TG; STA 449 F3 Jupiter, NETZSCH)) analysis were performed to examine the oxygen desorption properties and weight variations. The iodometric titration method was carried out to determine the oxygen nonstoichiometry at room temperature, while oxygen nonstoichiometry at elevated temperatures was calculated based on the weight loss recorded by TG analysis.29 The morphology of the fabricated symmetrical cell and anode-supported fuel cell was observed by field emission scanning electron microscopy (FE-SEM; HITACHI-S4800). The electrical conductivity in air, N2, and CO2 was studied via a four-probe DC method within the temperature range from 900 to 300 °C. The chemical stability toward CO2 poisoning was evaluated based on the XRD patterns and electrical conductivity. The electrical conductivity relaxation (ECR) was applied to obtain the coefficients of chemical bulk diffusion (Dchem) and chemical surface exchange (κchem), and a detailed procedure was described in our previous work.30 2.4. Electrochemical Performance. EIS of the symmetrical cell was obtained using a Solartron 1287 and 1260A electrochemical workstation. The symmetrical cell was measured between 550 and 800 °C in air at a flow rate of 100 mL min−1 (standard temperature and pressure, STP). A frequency ranging from 0.1 to 105 Hz and 10 mV amplitude were applied to the EIS measurement. CO2 tolerance and long-term stability were conducted at 600 °C in air for 60 h and were subsequently measured in 1 vol % CO2-containing atmosphere for another 60 h. The current−voltage (I−V) curve of an anodesupported fuel cell was measured using a Keithley 2420 digital source meter. During the measurement, hydrogen fuel was fed into the anode side at a flow rate of 80 mL min−1, while the cathode was exposed to ambient air. 2.5. Calculation of ABE. ABE, the enthalpy difference of formation of metal oxides, the sublimation of metal, and the oxygen dissociation (eq 1), was derived from the thermodynamic data (HSC 6.0). ABE =

⎞ 1 ⎛⎜ n ΔHΔmOn − mΔHA − DO2⎟ ⎠ CNA m ⎝ 2 ⎞ 1 ⎜⎛ n + ΔHBmOn − mΔHB − DO2⎟ ⎠ CNBm ⎝ 2

Figure 1. XRD patterns of SFHf and SF (a) and magnified XRD patterns at 2θ range from 46° to 48° (b) and from 32° to 34° (c), respectively, and Rietveld refinement of the XRD pattern of SFHf (d).

group. The corresponding lattice parameters of SFHf are calculated to be a = b = c = 3.895 Å, which are similar to those of SF-based perovskite-type oxides reported in our previous studies.22,23 TEM and high-resolution TEM (HRTEM) images combined with selected area electron diffraction (SAED) were performed to check the microstructure and crystallinity of SFHf. As shown in Figure 2a, SFHf particles with the particle size of larger than 0.5 μm are sintered together due to the high calcination temperature during the preparation. As indicated in Figure 2b,c, a high crystallinity of SFHf is achieved. Figure 2c shows wellcrystallized lattice planes with a fringe spacing of 0.274 nm, matching well with the distance between two (110) crystal

(1)

where CNA = 12 and CNB = 6 are the coordination numbers of A/B cations with O anion, ΔHAmOn and ΔHBmOn are the formation enthalpies of AmOn and BmOn, ΔHA and ΔHB represent the sublimation enthalpies of A and B metals, and DO2 represents the dissociation enthalpy of O2. For example, at room temperature, ΔHSr = 163.9 kJ mol−1, ΔHFe = 415.7 kJ mol−1, ΔHHf = 618.4 kJ mol−1, ΔHSrO = −591.5 kJ mol−1, ΔHFe2O3 = −820.5 kJ mol−1, and ΔHHfO2 = −1144.7 kJ mol−1 could be obtained from the thermodynamic database in HSC 6.0.

3. RESULTS AND DISCUSSION 3.1. Physical and Chemical Properties. The roomtemperature XRD pattern reveals that SFHf shows a wellcharacterized cubic structure after calcination at 1250 °C for 10 h (Figure 1). SF was also synthesized and characterized for comparison purpose. As shown in Figure 1b, SF contains several split peaks at ∼47°, but SF after doping with Hf4+ shows a reduced degree of the peak splitting, suggesting that Hf4+ doping could effectively stabilize the crystal structure. Figure 1c shows the magnified part near the main peak at ∼33°, where the peak of SFHf shifts to a lower diffraction angle when it is compared to the peak of SF, suggesting the lattice expansion after doping with Hf4+. This lattice expansion is ascribed to the replacement of partial Fe cations with Hf4+ since ionic radii are 0.645, 0.585, and 0.710 Å for Fe3+, Fe4+, and Hf4+ when they are assumed at the high-spin state and with 6-fold coordination. As shown in Figure 1d, Rietveld refinement of the XRD pattern of SFHf confirms a cubic crystal structure with Pm3̅m space

Figure 2. TEM image (a), magnified TEM image (b), HRTEM image (c), and selected area electron diffraction pattern (d) of SFHf. 2136

DOI: 10.1021/acsaem.8b00198 ACS Appl. Energy Mater. 2018, 1, 2134−2142

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homogeneous distribution without segregation of the corresponding Sr, Fe, and Hf elements in SFHf is observed. In order to investigate the effect of the doping level on the structure, SrFe0.8Hf0.2O3−δ with the doping level of 0.2 was also synthesized. Unfortunately, an undesired structure is formed (Supporting Information Figure S1), suggesting that a small amount of the doping level (e.g., 10 at. %) is critical for stabilizing the structure. Therefore, this work will only focus on SrFe0.9Hf0.1O3−δ. The phase structure stability of SFHf at elevated temperatures was also evaluated since the ionic radius and oxygen vacancy concentration in the lattice may dramatically change at operation temperature of 500−800 °C due to the thermal reduction of the cations. Figure 4a shows XRD patterns of SFHf annealed and then quenched from 500, 600, 700, and 800 °C to room temperature, respectively. For comparison purpose, unquenched SFHf is also included in Figure 4a. No extra or missing diffraction peaks are observed for the quenched SFHf compared with the as-synthesized SFHf, which excludes the phase transition of SFHf at elevated temperatures. Only the peak shift appears after quenching, where a higher quenching temperature leads to a slightly lower angle for the diffraction peak, suggesting the thermal reduction of the cations/the lattice expansion at elevated temperature. The structure stability of perovskite oxides after effective doping of redox-inactive cations in the B-site is in line with previous studies.31,32 We also evaluated the structure stability of SFHf at 600 and 800 °C in air for a prolonged time (120 h). As shown in Figure 4b, no impure phases are observed, suggesting that SFHf is a promising cathode for long-term SOFCs operation, while BSCF, a well-known perovskite oxide with a cubic structure, was reported to experience structural evolution from a cubic structure to a hexagonal structure after durable evaluation,33 likely due to the A-site cations mismatch (i.e., Ba2+ and Sr2+)

planes for a cubic perovskite structure. Figure 2d displays the SAED pattern of SFHf along the [011̅] zone axis. According to the distance of (100) and (111) planes and the angle (54.7°) between them, a perfect agreement with a cubic perovskite structure of SFHf is obtained as well. TEM-EDS characterization was further carried out to check the elemental distribution after doping Hf4+ in SFHf. As shown in Figure 3,

Figure 3. TEM image of SFHf (a) and EDS mappings of Sr (b), Fe (c), and Hf (d) in SFHf.

Figure 4. XRD patterns of SFHf quenched from 500, 600, 700, and 800 °C to room temperature (a), sintered at 600 and 800 °C in air for 120 h (b), and XRD patterns of the composites of SDC and SFHf after sintered at 600, 800, and 900 °C in air for 5 h (c). 2137

DOI: 10.1021/acsaem.8b00198 ACS Appl. Energy Mater. 2018, 1, 2134−2142

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Figure 5. Nyquist plots of the SFHf cathode in a SFHf|SDC|SFHf symmetrical cell (a), and derived polarization ASRs (b) and the corresponding activation energy (c) of the SFHf and SF cathode in symmetrical cells; electrochemical performance of the SFHf cathode using an anode-supported fuel cell (d).

Fex+ with the redox-inactive Hf4+ in the B-site that alters the electron hopping behavior.29,38 Nevertheless, the obtained electronic conductivity of SFHf is still acceptable and comparable to the state-of-the-art cathode materials such as BSCF.12 3.2. ORR Activity. We then carried out a series of characterizations to evaluate the ORR activity of SFHf as cathode for SOFCs. We determined the ORR activity of SFHf in an electrode|SDC|electrode symmetrical cell configuration. The corresponding morphology is shown in Figure S3, SFHf electrode with porous structures adheres firmly to the dense SDC electrolyte, ensuring sufficient active sites for ORR, low interfacial resistance, and stable electrochemical performance. SF electrode shows a similar morphology, indicating that the difference of ORR activity and stability between SFHf and SF is not from extrinsic factors; e.g., microstructure. Nyquist plots of a SFHf|SDC|SFHf symmetrical cell between 550 and 800 °C under open circuit voltage are shown in Figure 5a. For comparison purpose, Nyquist plots of a SF|SDC|SF symmetrical cell are shown in Figure S4. The intercept in the Nyquist plot at a relatively high frequency is the ohmic resistance mainly from the electrolyte resistance, as well as the electrode and connections. A small difference in ohmic resistance between SF and SFHf could be due to a small difference in the thickness of the electrolyte of symmetrical cells. The polarization ASRs are calculated from the intercept difference between the high and low frequency on the real axis in the Nyquist plot. A low ASR is an indicator of realizing the high ORR activity. As summarized in Figure 5b, the SFHf cathode delivers a high ORR activity with ASRs as low as 0.600, 0.193, 0.075, 0.037, 0.020, and 0.012 Ω cm2 at 550, 600, 650, 700, 750, and 800 °C, respectively, while for the SF electrode, ASRs of 1.011, 0.304, 0.113, 0.051, 0.030, and 0.014 Ω cm2 are recorded at 550, 600, 650, 700, 750, and 800 °C, respectively. The corresponding activation energy (Ea) is presented in Figure

and B-site Co cation in BSCF. Before evaluating the ORR activity and stability, we should ensure the compatibility between the electrolyte and the cathode. The compatibility is of importance since the SFHf electrode and the SDC electrolyte need to be sintered together at a high temperature (900 °C) to ensure a low interfacial resistance and a stable cell structure. We examined the compatibility by simply mixing SFHf and SDC powders with 1:1 weight ratio and then firing at 900 °C for 5 h, which is the same temperature for the fabrication of the SFHf cathode layer. For comparison purpose, the mixture without being fired and fired at 600 and 800 °C for 5 h was also examined. According to the XRD patterns shown in Figure 4c, no visible change to SFHf is observed, indicating a good chemical compatibility between SFHf and SDC. We also evaluated the electrical conductivity of SFHf since sufficiently high electronic conductivity is in general a basic requirement for obtaining high electrochemical performance, while insufficient electronic conductivity may probably lead to a large resistance. It was reported that the introduction of redoxinactive dopants decreased the electronic conductivity of perovskite oxides.29 Therefore, the effect of Hf4+ doping on the electronic conductivity was investigated. Figure S2 shows the temperature-dependent electronic conductivity of SF and SFHf within the temperature range of 300−900 °C in air. The electronic conductivity of both SFHf and SF increases with temperature until reaching a maximum and then decreases with the temperature, suggesting the change from a semiconducting behavior to a metallic conducting behavior. The transition point shifts to a lower temperature after Hf4+ doping, indicating that the reduction of Fe cations becomes easier,29,34−37 in line with TG and O2-TPD results (shown in section 3.2). In comparison to the undoped SF oxide, the electronic conductivity of SFHf is lower in the whole temperature range in air with the maximum electronic conductivity of ∼50 S cm−1. The decreased conductivity may arise from the substitution of 2138

DOI: 10.1021/acsaem.8b00198 ACS Appl. Energy Mater. 2018, 1, 2134−2142

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Figure 6. Room-temperature (a) and temperature-dependent oxygen nonstoichiometry (b) of SFHf and SF; electrical conductivity relaxation response curves of SFHf (c) and SF (d).

5c, where SFHf exhibits a relatively low Ea of 115 kJ mol−1 and SF shows Ea of 124 kJ mol−1. The obtained Ea is comparable to the state-of-the-art BSCF (116 kJ mol−1),11 indicating relatively high ORR activity maintains at intermediate-to-low temperature. These results clearly indicate that Hf4+ doping could effectively enhance the ORR activity. Figure S5 shows a comparison of ASRs between SFHf and other Co-free cathode materials based on Arrhenius form. For example, the electrochemical performance of SFHf cathode outperforms or at least is comparable to the other Co-free cathode materials such as Ba0.95Ca0.05Fe0.95In0.05O3−δ (0.038 Ω cm2 at 700 °C),39 Ba 0.95 La 0.05 FeO 3−δ (0.211 Ω cm 2 at 600 °C), 40 BaFe0.95Sn0.05O3−δ (0.21 Ω cm2 at 600 °C),41 and SmBaFe2O5+δ (1.348 Ω cm2 at 600 °C).42 Then electrochemical performance of the SFHf cathode was further evaluated using an anode-supported fuel cell (Ni+YSZ| YSZ, ∼5 μm|SDC, ∼2 μm|SFHf, ∼20 μm). The corresponding morphology is shown in Figure S6. As shown in Figure 5d, a typical potential and current density curve (I−V) and the corresponding current density and power density curve (I−P) were recorded. Peak power densities of 1935, 1417, 826, 528, and 275 mW cm−2 are achieved at the operation temperature of 750, 700, 650, 600, and 550 °C, respectively, suggesting superior electrochemical performance of the SFHf cathode. We note that these are initial results, which could be further improved. We further performed more characterizations to gain more insight into the improvement of the ORR activity. The thermal reduction is typically accompanied by the creation of oxygen vacancy, thus improving electrochemical performance. Therefore, O2-TPD and TG profiles were recorded to obtain the thermal reduction of reducible cations such as Fe cations in the lattice at elevated temperatures. Oxygen release and mass loss in TPD and TG profiles are clearly observed (Supporting Information Figure S7). Oxygen release accompanying the creation of the oxygen vacancy occurs at ∼375 and ∼420 °C for

SFHf and SF, and mass loss starts at ∼430 and ∼465 °C for SFHf and SF, respectively. Both characterizations suggest that the partial reduction of Fex+ cations in SFHf is easily realized, indicating that the Hf4+ doping facilitates the ionic mobility in the lattice of perovskite oxides. As shown in Figure S7a, there seem to be two main steps in the oxygen release process. In the first step, the α peak of SFHf is much larger than that of SF, suggesting the kinetics of oxygen release of SFHf is much faster than that of SF, while, in the second step indicated by β peak, the kinetics of oxygen release of SFHf gets slower than that of SF. TG profiles (Supporting Information Figure S7b) show similar change behavior: at first, the mass of SFHf decreases much more quickly than that of SF, and then the mass of SF decreases faster than that of SFHf at temperatures higher than ∼655 °C, which is in good agreement with the O2-TPD result. Therefore, the improved ORR activity of SFHf is ascribed to the facilitation of ionic mobility and the enhancement of oxygen desorption kinetics. To obtain quantitative data, temperature-dependent oxygen nonstoichiometry using iodometric titration and TG analysis was also measured. As shown in Figure 6a, SFHf shows a relatively low oxygen nonstoichiometry of ∼0.124 at room temperature, which is much smaller than that of the well-known cathode materials such as BSCF and SNC,20,25 suggesting that Hf4+ doping seems to be detrimental for the creation of oxygen vacancy at room temperature. We then further analyzed the oxygen nonstoichiometry at elevated temperatures. As shown in Figure 6b, the oxygen nonstoichiometry of SFHf at SOFCs operation temperature range is slightly over that of SF, ensuring relatively high ORR activity of SFHf. It is generally believed that oxygen vacancy plays a significant role in determining ORR activity since both oxygen surface exchange and bulk diffusion coefficients are closely related to oxygen vacancy.10,11,30 ECR technique shows that SFHf possesses much faster oxygen exchange and diffusion kinetics than SF (Figure 6c,d). For example, Dchem and κchem of SFHf are 4.96 × 10−5 cm2 s−1 and 2139

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Figure 7. Time-dependent polarization ASRs of SFHf at 600 °C in air and 1 vol % CO2, respectively, and (insets) corresponding Nyquist plots (a) and Nyquist plots of a SFHf|SDC|SFHf symmetrical cell at 600 °C in air and 1 vol % CO2, respectively (b). Ohmic resistances in all spectra are subtracted in order to allow a direct comparison of the polarization resistances.

4.12 × 10−4 cm s−1, respectively, at 700 °C (Supporting Information Figure S8), which are much higher than those of SF and comparable to the state-of-the-art BSCF30 and typical Co-free cathode materials (e.g., Bi0.5Sr0.5FeO3−δ).43 The corresponding activation energies of both Dchem and κchem of SFHf are also much lower than those of SF (Supporting Information Figure S8). 3.3. CO2 Tolerance and Long-Term Electrochemical Stability. The chemical stability (CO2 tolerance) and longterm electrochemical stability of the SFHf electrode were also evaluated. As shown in Figure 7a, polarization ASRs of the SFHf cathode obtained from a SFHf|SDC|SFHf symmetrical cell are considerably stable during the first 60 h operation in pure air at 600 °C. To further determine the CO2 resistance of SFHf, CO2 was then introduced into the cathode side. It should be noted that the polarization ASR increases suddenly from ∼0.193 to ∼0.356 Ω cm2 after switching the atmosphere to the presence of 1 vol % CO2, but the increased scale is still greatly smaller than those of available data for BSCF, SSNC, and SrNb0.1Co0.9−xFexO3−δ.20,25 For example, the polarization ASR of the SSNC cathode increases from ∼0.077 to ∼0.77 Ω cm2 (∼10 times). Importantly, the polarization ASR of SFHf negligibly increases during another 60 h operation, suggesting that only CO2 adsorption without structural destruction on the perovskite surface happens. In addition, the polarization ASR rapidly recovers to the original value after cutting the flow of CO2 (Figure 7b), suggesting that the CO2 adsorption on occupying active sites for ORR is reversible. As to the undoped SF electrode, the polarization ASR quickly increases from ∼0.481 (10 min) to ∼0.513 Ω cm2 (60 min) in 1 vol % CO2 (Supporting Information Figure S9), indicating the enhanced CO2 tolerance after Hf4+ doping. It has been widely recognized that chemically instable perovskite-type compositions such as BSCF showed substantial changes in XRD patterns after CO2 poisoning.20,21,44 Therefore, other than the electrochemical characterizations, powder XRD was also performed to probe CO2 poisoning. Figure S10a shows that the original cubic crystal structure of SFHf is retained without carbonate or impure phases observed after CO2 poisoning, while SF undergoes a clear phase change with extra XRD peaks observed. We also note that the electrical conductivity is a good indication for the evaluation of the chemical stability. As shown in Figure S10b, the conductivity of SFHf is similar in both N2 and CO2, while the conductivity of SF decreases dramatically in CO2 atmosphere. The above results successfully suggest that Hf doping obviously improves

the chemical stability and SFHf is a promising CO2-tolerant cathode for SOFCs operated in ambient air. ABE could be a descriptor to predict the CO2 tolerance.45 As reported, the more negative the ABE, the higher tolerance of the perovskite oxide against CO2.24,45 As shown in Figure S11, we find that ABE values for SFHf and SF show a similar trend at both room temperature and elevated temperature (600 °C). ABE values become more negative after doping with Hf4+ when compared to undoped SF. We also find that SFHf shows comparable ABE to well-known CO2-tolerant perovskite-type oxides such as Nb-doped SF, further demonstrating that SFHf tolerates CO2 effectively. In addition, oxygen vacancy has also been considered as one of the determining factors during the carbonate formation on the perovskite surface.45 Meanwhile the positively charged oxygen vacancy may enrich A-site dopants (e.g., Sr and Ba) on the surface via electrostatic attraction, thus easily forming carbonates on the surface.13 Therefore, a moderate oxygen nonstoichiometry is preferable in achieving stability. Based on our results, we might propose that roomtemperature and elevated-temperature oxygen nonstoichiometry of lower than 0.12 and 0.40 of perovskites is still workable for the CO2 tolerance.

4. CONCLUSIONS In this work, a Co-free CO2-tolerant Hf-doped SFHf perovskite oxide has been developed as a cathode for SOFCs. We find that SFHf after slightly doping Hf4+ exhibits high and stable ORR activity, and chemical stability. SFHf cathode has a polarization ASR as low as 0.193 Ω cm2 at 600 °C in a symmetrical cell and has a peak power density as high as 1.94 W cm−2 at 700 °C in an anode-supported fuel cell. The high ORR activity is related to the moderate oxygen vacancy concentration and electronic conductivity, and the fast oxygen kinetics at operation temperature. The improved chemical stability is attributed to the doping of the redox-inactive Hf cation in the Fe site of SF. This work suggests a facile yet effective strategy in the design of highly active and stable cathodes for SOFCs.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsaem.8b00198. Additional figures of XRD patterns, electronic conductivity, SEM, O2-TPD and TG, Dchem and κchem, ABE results, Nyquist plots, a comparison of ASRs, and timedependent polarization ASRs (PDF) 2140

DOI: 10.1021/acsaem.8b00198 ACS Appl. Energy Mater. 2018, 1, 2134−2142

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ACS Applied Energy Materials

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +86-20-85220223. Fax: +86-20-85220223. ORCID

Dengjie Chen: 0000-0002-3234-7020 Zongping Shao: 0000-0002-4538-4218 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by the National Nature Science Foundation of China (Grant No. 51702125), Fundamental Research Funds for the Central Universities (Grant No. 21616301), Pearl River S&T Nova Program of Guangzhou (Grant No. 201806010054), and the China Postdoctoral Science Foundation (Grant No. 2017M620401).

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