- Email: [email protected]
Calcium-doped ceria materials for anode of solid oxide fuel cells running on methane fuel
Journal of Power Sources 347 (2017) 79e85
Contents lists available at ScienceDirect
Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Calcium-doped ceria materials for anode of solid oxide fuel cells running on methane fuel Kai Zhao a, b, Yanhai Du a, * a b
College of Applied Engineering, Sustainability and Technology, Kent State University, Kent, OH 44242, United States The Voiland School of Chemical Engineering & Bioengineering, Washington State University, Pullman, WA 99164, United States
h i g h l i g h t s Ca doped ceria is found to be a potential anode component for hydrocarbon SOFC. Ce0.9Ca0.1O2-d shows similar conductivity compared with Sm or Gd doped ceria. Ni-Ce0.9Ca0.1O2-d is successfully applied as the anode of SOFC. Good performance stability is obtained in methane fuel.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 October 2016 Received in revised form 23 January 2017 Accepted 26 January 2017
A calcium-doped ceria with nominal compositions of Ce1-xCaxO2-d (0.00 x 0.30) has been developed as an anode component for solid oxide fuel cells running on methane fuel. Crystal phases of Ce1-xCaxO2d are investigated with respect to the amount of calcium dopant. The Ce1-xCaxO2-d shows single fluorite phase when the calcium is within 15 mol.%, and higher calcium doping levels lead to the appearance of a secondary phase (CaO). Conductivities of Ce1-xCaxO2-d ceramics are studied by a four-probe method in air and the composition of Ce0.9Ca0.1O2-d (x ¼ 0.10) is found exhibiting the highest conductivity among the samples investigated in this work. Electrocatalytic properties of Ce0.9Ca0.1O2-d are evaluated based on NiCe1-xCaxO2-d anode supported single cell running on methane fuel. At 800 C, the single cell with NiCe0.9Ca0.1O2-d (x ¼ 0.10) anode exhibits an optimum maximum powder density (618 mW cm2) and good performance stability during 30 h operation in methane fuel. The promising findings substantiate the good performance of Ni-Ce0.9Ca0.1O2-d anode for electrochemical oxidation of methane fuel. © 2017 Elsevier B.V. All rights reserved.
Keywords: Calcium-doped ceria ceramic Conductivity Composite anode Solid oxide fuel cell Methane fuel
1. Introduction Direct hydrocarbon solid oxide fuel cells (SOFCs) have attracted growing interest due to the wide availability of natural gas. The nickel-containing cermets (e.g., Ni-yttira-stabilized zirconia and Ni-samarium-doped ceria) have been regarded as typical anode components for the high conductivity and good electrochemical activity [1]. The nickel, however, catalyzes the decomposition of hydrocarbon fuels, resulting in severe deposition of carbon in the anode [2,3]. The growth of carbon in the anode could deteriorate the microstructure and induce the coking problem. These phenomena would lead to the degradation of electrocatalytic properties of the anode and loss of electrochemical performance for the
* Corresponding author. E-mail address: [email protected] (Y. Du). http://dx.doi.org/10.1016/j.jpowsour.2017.01.113 0378-7753/© 2017 Elsevier B.V. All rights reserved.
fuel cells. Therefore, suppressing carbon deposition on the surface of Ni-containing anode appears as an important issue to be solved. Ceria-based materials have received some attention as a reforming catalyst in automobile since it can release and uptake oxygen by the following chemical reaction [4,5]:
CeO2 4CeO2x þ ðx=2ÞO2
(1)
The oxygen storage capacity characteristic could be used to catalyze the electrochemical oxidation of hydrocarbon fuels and suppress the carbon deposition in the anode. The catalytic properties of ceria have been evaluated based on SOFC single cells running on methane fuel [6]. The single cell with ceria catalyst shows evidently improved performance in the initial operation for 14 h in methane fuel. Carbon deposition phenomenon, however, can still be identified in the following 6 h measurement. In order to enhance performance stability of single cells in hydrocarbon fuels,
80
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
research effort has been devoted toward optimization of the catalytic properties for ceria-based materials. Recent research results indicate that catalytic properties of ceria are essentially related to their crystal structures and defect properties [7e9]. Various elements, such as lanthanum, zirconium and calcium, have been doped into ceria to enhance the electrocatalytic property. Among them, the calcium-doped ceria shows highest electronic conductivity and oxygen ionic mobility, which are beneficial for electrochemical oxidation of hydrocarbon fuels. The doping of calcium into CeO2 produces lattice oxygen vacancy by a charge-compensating mechanism [10]:
ð1 xÞCeO2 þ xCaO/Ce1x Cax O2x þ zVO,,
(2)
The introduction of calcium increases the oxygen transport properties of ceria through the creation of the oxygen vacancy defects, favoring the conduction of oxygen ions. The increased oxygen ion mobility could facilitate the oxidation of hydrocarbon fuels, suppressing the carbon deposition in the anode. In addition, the doping of calcium will increase the lattice parameters of ceria, thus it has been used as a promoter for oxygen storage capacity of ceria. Compared to undoped ceria, the calcium-doped ceria demonstrates a higher oxygen storage capacity (~40% higher than CeO2), and it displays a remarkably increased redox activity [11,12]. Due to these promising results, it has been used as a catalyst for reforming of hydrocarbon fuels [13,14]. As calcium-doped ceria exhibits both good oxygen ionic conductivity and some catalytic property for hydrocarbon fuel reforming, applying the material in the anode of SOFC would make the best use of its electrochemical function and catalytic characteristic. In this case, different oxygen sources (e.g., oxygen ions from the cathode, and oxygen from steam) in the anode can be effectively utilized to oxidize carbon species and suppress the carbon deposition problem. On the other aspect, the oxygen vacancy can be trapped by the defect clusters at high levels of element doping, lowering the conductivity of oxygen ions [15]. Thus further research is necessary to optimize the material composition for improving the electrochemical property. In this work, Ce1-xCaxO2-d (0 x 0.30) powders are synthesized by a urea combustion method. The crystal structure and conductivities of Ce1-xCaxO2-d are investigated with respect to the amount of calcium content. Catalytic properties of Ce1-xCaxO2-d are evaluated based on Ni-Ce1-xCaxO2-d anode supported single cells running on methane fuel. The optimum calcium doping amount in Ce1-xCaxO2-d is determined with respect to conductivity of Ce1xCaxO2-d ceramics, electrochemical performance and performance stability of the single cells running on methane fuel. 2. Material and methods 2.1. Powder preparation The Ce1-xCaxO2-d (0 x 0.30) powders were synthesized by a urea combustion method. Reagent grade Ca(NO3)2$4H2O, Ce(NO3)3$6H2O and CO(NH2)2 (99.99% purity from Sigma-Aldrich) were used as starting chemicals. The nitrates were weighed according to the nominal compositions and dissolved into deionized water to form a solution. CO(NH2)2 was added into the solution as a combustion agent. The mole ratio of CO(NH2)2 to the total metal cation content was determined to be two [16]. The solution was stirred for 1 h and heated on a hot plate until auto-ignition and self-sustaining combustion occurred. The obtained ash was subsequently grounded for 24 h in ethanol medium using a ball milling process with zirconia balls. Then the powder was calcined at 900 C for 2 h in air. Ce0.8Sm0.2O1.9 (SDC) electrolyte powder and NiO-Ce1-xCaxO2d anode powder with 60 wt% NiO and 40 wt% Ce1-xCaxO2-d were
synthesized by the same urea combustion process. La0.6Sr0.4Co0.2Fe0.8O3-d (LSCF) cathode powder was synthesized by a glycinenitrate combustion method. The detailed synthetic procedures have been published previously [16,17]. 2.2. Single cell fabrication Anode supported single cell with a configuration of porous NiOCe1-xCaxO2-d anode supporter/yttira-stabilized zirconia (YSZ)/SDC bi-layer electrolyte/LSCF cathode was fabricated by a dry-pressing and spin-coating process. The NiO-Ce1-xCaxO2-d was mixed with a pore former (corn starch, Sigma-Aldrich) at 15 wt% by a ball-milling process. The mixed powder was compressed into disc pellets at a pressure of 80 MP using a hydraulic press machine. The diameter of green pellets was 21 mm and the thickness was ~2.5 mm. The pellets were calcined at 1000 C for 3 h in air to obtain sufficient mechanical strength for the following spin-coating processes. 17 wt% YSZ powder (Tosoh, Japan) was mixed with a lab-made organic carrier to form a slurry for the spin-coating process. The organic carrier was composed of B73210 organic binder (30 wt%, Ferro Electronics Materials), terpineol (40 wt%, Sigma-Aldrich) and ethanol (30 wt%, Decon Labs Inc). The slurry was ball-milled for 24 h and spin-coated on the surface of the NiO-Ce1-xCaxO2-d substrate at a rotation speed of 4500 rpm for 30 s. The spin-coating process was repeated for 3 times in order to achieve a ~10 mm YSZ layer. To avoid the formation of pin holes and microcracks in the electrolyte, the layer was calcined at 500 C in air for 30 min before the deposition of another layer. The SDC layer was fabricated on the surface of YSZ layer in the same way. The NiO-Ce1-xCaxO2d anode substrate and YSZ/SDC bi-layer electrolyte was co-sintered at 1400 C for 4 h in air. After the co-sintering process, the anode supported half cell showed a diameter of 17 mm and a thickness of ~2 mm. The LSCF cathode layer with an area of 0.79 cm2 was deposited on the surface of SDC electrolyte by the similar spin-coating process. 14.1 wt% LSCF powder was dispersed in an organic slurry consisting of 85.1 wt% ethanol and 0.8 wt% ethyl cellulose (SigmaAldrich) by the ball-milling process. The cathode slurry was spincoated on the surface of SDC electrolyte at a rotation speed of 3000 rpm. Then, the cathode layer was sintered at 1000 C for 2 h in air. 2.3. Structural characterization Crystal phases of the Ce1-xCaxO2-d powders and ceramics were studied by the X-ray diffractometry (XRD, Rigaku Smart-Lab XRD) using Cu Ka radiation at 40 kV and 44 mA. The data were collected in the continuous mode in the 2q range of 20e80 at 0.5 min1. Microstructures of the sintered Ce1-xCaxO2-d ceramics and the NiCe1-xCaxO2-d anode supported single cells were investigated by a scanning electron microscope (SEM, FEI Quanta 450). 2.4. Conductivity measurement Conductivities of Ce1-xCaxO2-d ceramics were measured by direct current (DC) four-probe technique and alternating current (AC) impedance measurement (Metrohm Autolab 302N Electrochemical workstation with FRA32 Impedance Analyzer), respectively. For the DC four-probe measurement, a dense rectangular bar sample (length: 25.0 mm, width: 5.4 mm, and height: 4.0 mm) was prepared and four parallel silver-line electrodes were used in the DC conductivity test. In the case of AC impedance measurement, a circular disc sample (diameter: 8.8 mm and thickness: 1.0 mm) was used with both sides painted with silver paste. The test was carried
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
out at the input sinuous voltage of 20 mV in the frequency range of 1 Hze10 MHz.
81
d powder is investigated by image analysis method using Image
Electrochemical performance of the single cell with the Ni-Ce1anode was characterized using a lab-designed fuel cell testing system. The system consists of an Autolab electrochemical workstation (Metrohm 302N and 20 A current booster), a set of mass flow controllers (P4B, MKS Instruments) and alumina sample holders. An alumina-based ceramic binder (AREMCO Products) was used as the sealing material. The single cell was heated up to the testing temperature at 3 C min1 and a hydrogen flow rate of 100 ml min1 was used to reduce the nickel oxide in the anode, while the cathode side was exposed to air. After the open-circuit voltage (OCV) of the single cell reached a stable state, electrochemical performance of the single cell was tested in hydrogen as fuel. Then, the anode fuel was switched from hydrogen to humidified methane. A methane flow rate of 100 ml min1 was fed into a water bubbling system at the room temperature (25 C) to make humidified methane. The humidified methane was delivered to the anode of the single cell to study the electrochemical performance. Electrochemical impedance spectra of the single cell were obtained under the OCV condition in the frequency range of 0.01 Hze100 kHz. The amplitude of the input sinuous signal was set to be 10 mV for the impedance measurement. Performance stability of the single cell was evaluated at 800 C in humidified methane fuel as well. The methane flow rate was set to be 100 ml min1 in the anode, and the cathode was exposed to air. The performance stability was tested by monitoring voltages of the single cell at a constant current load of 500 mA cm2.
pro-plus 6.0 software [19]. For each sample, three SEM images from different parts of the powder are used to estimate the particle size distribution and the average particle size is determined to be ~300 nm by Guass fitting. The Ce1-xCaxO2-d powders with different calcium contents show similar morphology and particle size distribution in this work. Fig. 3 (a) shows XRD data of Ce1-xCaxO2-d (0 x 0.15) ceramics sintered at 1400 C. The ceramics exhibit single cubic fluorite phase by the XRD analyses. Fig. 3 (b) shows the short range (28e34 ) XRD spectra data. With the increase of calcium doping amount, diffraction peaks corresponding to (111) and (200) planes of the cubic fluorite structure shift to lower 2q ranges. The Rietveld refinement technique has been applied to evaluate the crystallographic parameters of the cubic Ce1-xCaxO2-d ceramics. The lattice parameters are listed in Fig. 3 (c). The lattice parameters generally increase with calcium content in Ce1-xCaxO2-d. This phenomenon can be ascribed to the larger ion radius of Ca2þ (1.14 Å) compared with that of Ce4þ (1.01 Å) [9]. Moreover, the variation of the lattice parameter basically obeys the Vegard's rule in this work [20]. Fig. 4 shows SEM images of Ce1-xCaxO2-d (0 x 0.15) ceramics sintered at 1400 C in air. The samples exhibit generally dense microstructure with few isolated pores. Archimedes' method has been used to estimate the density of the ceramics and the relative densities are found to be over 95% for all the samples under this investigation [21]. Moreover, the relative density displays an increasing trend with the calcium amount (shown in Supplementary Material Fig. S1), which is in agreement with the reduced pores observed in the Ce1-xCaxO2-d ceramics. Thus, for the calcium-doped ceria compounds, the calcium acts as a sintering agent [22].
3. Results and discussion
3.2. Conductivity
3.1. Structure characterization
AC impedance measurement has been carried out using Ce1(x ¼ 0.00, 0.05, 0.10 and 0.15) ceramic disc samples sintered at 1400 C. Fig. 5 shows Nyquist plots of the Ce1-xCaxO2-d ceramics obtained at 300 C in air. The Nyquist plots are characterized by high frequency semi-circles and a low frequency line feature. To distinguish resistances from grains and grain boundaries of Ce1CaxO2-d ceramics, the impedance spectra are fitted by an equivax lent circuit model ðR1 Q1 ÞðR2 Q2 ÞðR3 Q3 Þ (shown in the insert of Fig. 5(a)). R is the resistance and Q represents the constant phase element. The impedance ZQ of a constant phase element Q and the equivalent capacitance C can be calculated according to the following equations [23,24]:
2.5. Cell performance measurement
xCaxO2-d
Fig. 1 shows XRD patterns of Ce1-xCaxO2-d powders calcined at 900 C for 2 h in air. The characteristic diffraction peaks for the fluorite phase are identified with the calcium doping amount in the range of 0e30 mol.% (0 x 0.30). At high calcium doping levels (x ¼ 0.20 and 0.30), an additional small peak is observed, which is attributed to the CaO phase in the composite [11,18]. The results indicate a solubility limit of 15 mol.% for calcium-doped ceria materials, which is in agreement with previous published results [11]. Fig. 2 (a) shows SEM image of the Ce0.9Ca0.1O2-d powder calcined at 900 C in air. The powder consists of fine particles with some agglomerations. Particle size distribution of the Ce0.9Ca0.1O2-
xCaxO2-d
ZQ ¼ 1
Q ði$2pf Þn
. C ¼ ðR$Q Þ1=n R
Fig. 1. XRD patterns of Ce1-xCaxO2-d (0 x 0.30) powders calcined at 900 C in air.
(3) (4)
where i is the imaginary unit, f is the frequency, and n is the an exponent parameter. Table 1 lists the fitting results obtained from AC impedance spectra. Generally, the equivalent capacitance C1 for ðR1 Q1 Þ is on the orders of magnitude of 1011 F, thus, R1 can be attributed to the resistance from grains of Ce1-xCaxO2-d ceramics. While, the values of C2 are on the orders of magnitude of 108 F. Therefore, R2 can be regarded to be the grain boundary resistance. The equivalent capacitances from ðR3 Q3 Þ component vary in the range of 106-104 F, the large capacitances are attributed to the electrode processes on both surfaces of Ce1-xCaxO2-d samples [22,25]. The sample without Ca doping shows one discernable semicircle, and the total resistance (sum of grain and grain boundary
82
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
Fig. 2. (a) SEM image of Ce0.9Ca0.1O2-d powder calcined at 900 C in air and (b) Particle size distribution of the Ce0.9Ca0.1O2-d powder.
Fig. 3. XRD patterns of Ce1-xCaxO2-d (0 x 0.15) ceramics sintered at 1400 C in air: (a) 2q range of 20e80 and (b) 2q range of 28e34 . (c) Lattice parameters obtained from XRD analyses.
resistances) is determined to be 45,129 U by impedance fitting (shown in Fig. 5 (a)). With the doping of calcium, grain and grain boundary resistances can be resolved from the impedance spectra (shown in the Nyquist plots in Fig. 5 and in the impedance fitting results in Table 1). For all the samples under the investigation, the gain resistance accounts as major contribution to the total resistance of the ceramics. Thus, the variation of the conductivity (or resistance) can be ascribed to the effects of calcium doping on the ceria-based materials. When increasing the calcium doping amount from 0 mol.% to 10 mol.%, the grain resistance displays a significant reduction. On the one hand, partially replacing tetravalent cerium with divalent calcium could give rise to the formation of oxygen vacancies: CeO2
00
CaO!CaCe þ VO,, þ OxO
(5)
The increased number of charge carriers improves the conductivities of the ceramics and reduces the grain resistances. On the
other hand, the introduction of calcium leads to the expansion of crystal structure (Fig. 3 (c)), enhancing mobility of the oxygen ions [13,22]. The contribution factors from these two aspects could lower resistances of the ceramics. Additionally, the grain boundary resistance exhibits a decreasing trend with calcium doping as well. This could be attributed to the increased relative density of the ceramics (shown in Supplementary Material Fig. S1). Calcium has been reported to be a sintering promoter for ceria-based materials [22]. The increment of relative density could facilitate the transport of charge carriers at grain boundaries, and enhance total conductivities of the ceramics. At a higher calcium doping amount (e.g., 15 mol.%), the grain resistance displays an increasing trend (from 6054 to 6453 U at 300 C). This phenomenon could be associated with the formation of defect clusters, resulting in the loss of movable oxygen vacancies in the ceramic [14,15]. As a result, despite the reduced resistance from grain boundaries, the total resistance of the composite
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
83
Fig. 4. SEM images of Ce1-xCaxO2-d ceramics sintered at 1400 C: (a) x ¼ 0.00, (b) x ¼ 0.05, (c) x ¼ 0.10 and (d) x ¼ 0.15.
Fig. 5. Nyquist plots of Ce1-xCaxO2-d ceramics at 300 C in air: (a) x ¼ 0.00, (b) x ¼ 0.05, (c) x ¼ 0.10 and (d) x ¼ 0.15. The equivalent circuit model is shown in the insert of (a).
Table 1 Resistances of Ce1-xCaxO2-d ceramics at 300 C obtained from impedance spectra fitting. Ce1-xCaxO2-d x x x x
¼ ¼ ¼ ¼
0.00 0.05 0.10 0.15
C1 (F) 1.64 2.45 1.87 1.91
1011 1011 1011 1011
R1 (U)
C2 (F)
R2 (U)
C3 (F)
45,129 12,054 6054 6453
e 3.73 108 2.96 108 2.76 108
e 2406 595 392
6.32 1.32 4.74 4.28
(Ce0.85Ca0.15O2-d) shows an increasing trend. Fig. 6 shows conductivities of calcium-doped ceria ceramics obtained by the four-probe method at 800 C in air. The data derived from AC impedance spectroscopy tests are also included in the figure. The undoped ceria presents fairly low conductivity (0.002 S cm1 at 800 C), which is in agreement with previous reported data on ceria materials [26]. With the introduction of calcium, the conductivity improves significantly and the 10 mol.%
R3 (U) 106 104 104 104
642,278 85,542 37,928 40,694
calcium-doped sample (Ce0.9Ca0.1O2-d) exhibits a highest conductivity (0.102 S cm1 at 800 C) among all the specimens investigated in this work. The conductivity is competitive compared with the widely used samarium-doped ceria (0.082 S cm1 at 800 C) and gadolinium-doped ceria (0.053 S cm1 at 750 C) materials [27,28]. The high conductivity is believed to be favorable for electrochemical oxidation of hydrocarbon fuels in the anode. In addition, calcium is a much cheaper and widely available material in
84
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
Fig. 6. Conductivity of Ce1-xCaxO2-d ceramics at 800 C in air.
Fig. 7. SEM image of Ni-Ce0.9Ca0.1O2-d anode supported single cell.
comparison with samarium and gadolinium. Application of the material could help reduce the cost of SOFCs. Thus, 5e15 mol.% calcium-doped ceria are applied into the anode of the single cell consisting of Ni-Ce1-xCaxO2-d anode supporter/YSZ/SDC bi-layer electrolyte/LSCF cathode. 3.3. Single cell performance Fig. 7 shows a typical cross-sectional SEM image of the single cell. Image analysis method is used to estimate thicknesses of cell functional layers. The LSCF cathode layer exhibits good adhesion to the SDC layer and the thickness is determined to be ~10 mm. The SDC buffer layer appears to be dense with some isolated pores. The YSZ electrolyte layer displays a dense microstructure and the thickness is ~9 mm. The Ni-Ce0.9Ca0.1O2-d anode substrate presents
highly porous microstructure, which could be due to the burning out of pore former in the cell fabrication process. Fig. 8 (a) shows voltage and power density of the single cell with Ni-Ce1-xCaxO2-d (x ¼ 0.05, 0.10 and 0.15) anodes at 800 C in humidified methane fuel. The single cells present similar OCV of ~1.02 V, indicating a dense electrolyte layer and good sealing for the single cell. The maximum power density exhibits a slight variation as a function of calcium content in the Ce1-xCaxO2-d. The NiCe0.9Ca0.1O2-d (x ¼ 0.10) anode supported single cell displays an optimum maximum power density of 618 mW cm2 at 800 C in methane fuel. The performance is comparable with the typical NiCe0.8Sm0.2O1.9 anode supported single cell (572 mW cm2 at 800 C in methane fuel, shown in Supplementary Material Fig. S2) under the same operation condition. Electrochemical impedance spectra of the single cells are obtained under the OCV condition in methane fuel. Fig. 8 (b) shows Nyquist plots of the single cell at 800 C. The impedance spectra are fitted by the LRohm ðRH QH ÞðRL QL Þ equivalent circuit model (shown in the insert of Fig. 8(b)). The inductance element L represents the inductive impedance response, while Rohm is the ohmic resistance of the entire cell. ðRH QH ÞðRL QL Þ components denote the electrochemical processes from both the anode and the cathode. The fitting results of the electrochemical impedance spectra are listed in Table 2. The equivalent capacitance for the higher frequency impedance arc is on the order of magnitude of 103-102 F cm2. The electrode processes can be ascribed to the charge transfer process ðRH QH Þ. The lower frequency arc has a larger equivalent capacitance (~100 F cm2), which can be attributed to the gas adsorption and dissociation, and gas phase diffusion processes ðRL QL Þ. The total polarization resistance of the whole cell ðRp Þ is the sum of RH and RL [29e31]. The single cells illustrate similar ohmic resistances (0.26e0.27 U cm2, at 800 C in Table 2) with different calcium contents in Ce1-xCaxO2-d. For the nickel containing anode supported single cell with YSZ as electrolyte, the ohmic resistance from the electrolyte can be regarded as a major contribution to the ohmic loss of the whole cell. The similar ohmic resistance indicates a comparable electrolyte layers for different cells in this research. As the fabrication process of the LSCF cathode layer is kept same in this work, the variation of the polarization resistance (Rp ) could be attributed to differences in the anode. By comparing variations of RH and RL in Table 2, the differences of the total polarization resistance (Rp ) are due to the charge transfer process (RH ). As the Ce0.9Ca0.1O2-d demonstrates a higher conductivity among all the samples investigated in this work (shown in Figs. 5 and 6), the good conductivity could enhance the transfer of oxygen between the anode and the fuel, thus, improving the charge transfer processes at the anode [13]. This data is in agreement with the higher maximum power density obtained for the cell using Ni-Ce0.9Ca0.1O2-d as the anode.
Fig. 8. (a) Cell voltage and power density of Ni-Ce1-xCaxO2-d (x ¼ 0.05, 0.10 and 0.15) anode supported single cells at 800 C in methane, (b) Nyquist plots of the single cells under OCV condition, and (c) the performance stability. The equivalent circuit model is shown in the insert of (b).
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
85
Table 2 Electrochemical parameters of the single cells at 800 C in methane. Ni-Ce1-xCaxO2-d anode
Rohm (U cm2)
CH (F cm2)
RH (U cm2)
CL (F cm2)
RL (U cm2)
Rp (U cm2)
x ¼ 0.05 x ¼ 0.10 x ¼ 0.15
0.27 0.26 0.27
1.51 102 9.83 103 1.29 102
0.20 0.12 0.17
1.81 2.12 1.65
0.12 0.10 0.11
0.32 0.22 0.28
Fig. 8 (c) shows performance stability of the Ni-Ce1-xCaxO2-
Appendix A. Supplementary data
d (x ¼ 0.05, 0.10 and 0.15) anode supported single cells in methane
fuel at 800 C. The performance stability is investigated under the galvanostatic mode (at the current density of 500 mA cm2). Among all the single cells investigated in this research, the NiCe0.9Ca0.1O2-d anode supported cell displays a lowest degradation rate (1.1 mV h1) during 30 h operation in methane fuel. The degradation rate is comparable with typical Ni-Ce0.8Sm0.2O1.9 anode supported single cell (~1.6 mV h1) in the initial 10 h operation (shown in Supplementary Material Fig. S2). However, the overall stability for the Ni-Ce0.9Ca0.1O2-d anode supported cell is evidently enhanced compared with the Ni-Ce0.8Sm0.2O1.9 supported cell. After the performance stability test, SEM/EDX has been used to estimate the amount of carbon in the anode. Only a small amount of carbon (4.8 wt%) is detected in the Ni-Ce0.9Ca0.1O2d anode, while 16 wt% carbon is found out in Ni-Ce0.8Sm0.2O1.9 anode (shown in Supplementary Material Fig. S3 and Table S1). The higher tolerance of Ni-Ce0.9Ca0.1O2-d anode to carbon deposition suggests the application potential of calcium-doped ceria material (Ce0.9Ca0.1O2-d) as an anode component for SOFCs running on methane fuel. 4. Conclusions Calcium-doped ceria with nominal compositions of Ce1-xCaxO2d (0.00 x 0.30) is synthesized by the urea-combustion method. Crystal phases of the Ce1-xCaxO2-d powders are investigated with
respect to the calcium doping amount. The powders show single fluorite phase when the doping level is within 15 mol.% (x 0.15), and higher calcium doping leads to the formation of secondary CaO phase. Conductivities of the Ce1-xCaxO2-d ceramics are studied as a function of calcium amount as well. The 10 mol.% calcium-doped sample (Ce0.9Ca0.1O2-d) exhibits a high conductivity of 0.102 S cm1 at 800 C, which is competitive as compared with typical samarium- or gadolinium-doped ceria. The calcium-doped ceria is applied in the Ni-Ce1-xCaxO2-d anode supported single cell, and the cell with the Ni-Ce0.9Ca0.1O2-d anode gives an optimum maximum powder density (618 mW cm2 at 800 C) and reasonable performance stability during the ~30 h test in methane fuel. The promising electrochemical results suggest the application feasibility of the Ce0.9Ca0.1O2-d materials as the anode component of SOFCs. Acknowledgments The authors acknowledge the support from Kent State University and NASA (Contract No.: NNX14CC37P and NNX15CC12C).
Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jpowsour.2017.01.113. References [1] D. Sarantaridis, A. Atkinson, Fuel Cells 7 (2007) 246e258. [2] M.S. Khan, S.B. Lee, R.H. Song, J.W. Lee, T.H. Lim, S.J. Park, Ceram. Int. 42 (2016) 35e48. [3] N. Mahato, A. Banerjee, A. Gupta, S. Omar, K. Balani, Prog. Mater. Sci. 72 (2015) 141e337. [4] A. Trovarelli, Cat. Rev. 38 (1996) 439e520. [5] M. Ozawa, M. Kimura, A. Isogai, J. Alloys Compd. 193 (1993) 73e75. [6] Y. Yin, S. Li, C. Xia, G. Meng, J. Power Sources 167 (2007) 90e93. [7] N. Jaiswal, S. Upadhyay, D. Kumar, O. Parkash, J. Power Sources 222 (2013) 230e236. [8] C. Zhang, X.D. Wen, B.T. Teng, Y. Zhao, M. Fan, Fuel Process. Technol. 131 (2015) 1e6. € fer, Acta Mater. 103 (2016) 361e369. [9] A. Maheshwari, H.D. Wiemho [10] Y. Zheng, Y. Shi, H. Gu, L. Gao, H. Chen, L. Guo, Mater. Res. Bull. 44 (2009) 1717e1721. [11] V. Thangadurai, P. Kopp, J. Power Sources 168 (2007) 178e183. [12] Y. Zhang, S. Andersson, M. Muhammed, Appl. Catal. B Environ. 6 (1995) 325e337. [13] H. Song, U.S. Ozkan, J. Phys. Chem. A 114 (2010) 3796e3801. jcik, [14] S. de Carolis, J.L. Pascual, L.G.M. Pettersson, M. Baudin, M. Wo K. Hermansson, A.E.C. Palmqvist, M. Muhammed, J. Phys. Chem. B 103 (1999) 7627e7636. [15] M. Mogensen, T. Lindegaard, U.R. Hansen, G. Mogensen, J. Electrochem. Soc. 141 (1994) 2122e2128. [16] Q. Xu, D.P. Huang, K. Zhao, W. Chen, M. Chen, B.H. Kim, Ceram. Int. 37 (2011) 913e920. [17] K. Zhao, Q. Xu, D.P. Huang, M. Chen, B.H. Kim, Ionics 17 (2011) 247e254. [18] M. Yan, T. Mori, J. Zou, F. Ye, D.R. Ou, J. Drennan, Acta Mater. 57 (2009) 722e731. [19] M. Rochoux, Y. Guo, Y. Schuurman, D. Farrusseng, Phys. Chem. Chem. Phys. 17 (2015) 1469e1481. [20] S. Zha, C. Xia, G. Meng, J. Power Sources 115 (2003) 44e48. [21] K.Y. Park, T.H. Lee, J.T. Kim, N. Lee, Y. Seo, S.J. Song, J.Y. Park, J. Alloys Compd. 585 (2014) 103e110. [22] S. Banerjee, P.S. Devi, D. Topwal, S. Mandal, K. Menon, Adv. Funct. Mater. 17 (2007) 2847e2854. [23] C. Zhang, S. Wang, R. Guo, L. Liu, Y. Yang, J. Alloys Compd. 646 (2015) 283e290. [24] M. Zunic, G. Brankovic, C.R. Foschini, M. Cilense, E. Longo, J.A. Varela, J. Alloys Compd. 563 (2013) 254e260. [25] K.C. Anjaneya, G.P. Nayaka, J. Manjanna, G. Govindaraj, K.N. Ganesha, J. Alloys Compd. 578 (2013) 53e59. [26] S.R. Bishop, T.S. Stefanik, H.L. Tuller, J. Mater. Res. 27 (2012) 2009e2016. [27] M. Chen, B.H. Kim, Q. Xu, B.K. Ahn, W.J. Kang, D.P. Huang, Ceram. Int. 35 (2009) 1335e1343. [28] A.Z. Liu, J.X. Wang, C.R. He, H. Miao, Y. Zhang, W.G. Wang, Ceram. Int. 39 (2013) 6229e6235. [29] C. Lalanne, F. Mauvy, E. Siebert, M.L. Fontaine, J.M. Bassat, F. Ansart, P. Stevens, J.C. Grenier, J. Eur. Ceram. Soc. 27 (2007) 4195e4198. [30] E.C. Shin, P.A. Ahn, H.H. Seo, J.M. Jo, S.D. Kim, S.K. Woo, J.H. Yu, J. Mizusaki, J.S. Lee, Solid State Ionics 232 (2013) 80e96. [31] V.C. Kournoutis, F. Tietz, S. Bebelis, Fuel Cells 9 (2009) 852e860.
Contents lists available at ScienceDirect
Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Calcium-doped ceria materials for anode of solid oxide fuel cells running on methane fuel Kai Zhao a, b, Yanhai Du a, * a b
College of Applied Engineering, Sustainability and Technology, Kent State University, Kent, OH 44242, United States The Voiland School of Chemical Engineering & Bioengineering, Washington State University, Pullman, WA 99164, United States
h i g h l i g h t s Ca doped ceria is found to be a potential anode component for hydrocarbon SOFC. Ce0.9Ca0.1O2-d shows similar conductivity compared with Sm or Gd doped ceria. Ni-Ce0.9Ca0.1O2-d is successfully applied as the anode of SOFC. Good performance stability is obtained in methane fuel.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 26 October 2016 Received in revised form 23 January 2017 Accepted 26 January 2017
A calcium-doped ceria with nominal compositions of Ce1-xCaxO2-d (0.00 x 0.30) has been developed as an anode component for solid oxide fuel cells running on methane fuel. Crystal phases of Ce1-xCaxO2d are investigated with respect to the amount of calcium dopant. The Ce1-xCaxO2-d shows single fluorite phase when the calcium is within 15 mol.%, and higher calcium doping levels lead to the appearance of a secondary phase (CaO). Conductivities of Ce1-xCaxO2-d ceramics are studied by a four-probe method in air and the composition of Ce0.9Ca0.1O2-d (x ¼ 0.10) is found exhibiting the highest conductivity among the samples investigated in this work. Electrocatalytic properties of Ce0.9Ca0.1O2-d are evaluated based on NiCe1-xCaxO2-d anode supported single cell running on methane fuel. At 800 C, the single cell with NiCe0.9Ca0.1O2-d (x ¼ 0.10) anode exhibits an optimum maximum powder density (618 mW cm2) and good performance stability during 30 h operation in methane fuel. The promising findings substantiate the good performance of Ni-Ce0.9Ca0.1O2-d anode for electrochemical oxidation of methane fuel. © 2017 Elsevier B.V. All rights reserved.
Keywords: Calcium-doped ceria ceramic Conductivity Composite anode Solid oxide fuel cell Methane fuel
1. Introduction Direct hydrocarbon solid oxide fuel cells (SOFCs) have attracted growing interest due to the wide availability of natural gas. The nickel-containing cermets (e.g., Ni-yttira-stabilized zirconia and Ni-samarium-doped ceria) have been regarded as typical anode components for the high conductivity and good electrochemical activity [1]. The nickel, however, catalyzes the decomposition of hydrocarbon fuels, resulting in severe deposition of carbon in the anode [2,3]. The growth of carbon in the anode could deteriorate the microstructure and induce the coking problem. These phenomena would lead to the degradation of electrocatalytic properties of the anode and loss of electrochemical performance for the
* Corresponding author. E-mail address: [email protected] (Y. Du). http://dx.doi.org/10.1016/j.jpowsour.2017.01.113 0378-7753/© 2017 Elsevier B.V. All rights reserved.
fuel cells. Therefore, suppressing carbon deposition on the surface of Ni-containing anode appears as an important issue to be solved. Ceria-based materials have received some attention as a reforming catalyst in automobile since it can release and uptake oxygen by the following chemical reaction [4,5]:
CeO2 4CeO2x þ ðx=2ÞO2
(1)
The oxygen storage capacity characteristic could be used to catalyze the electrochemical oxidation of hydrocarbon fuels and suppress the carbon deposition in the anode. The catalytic properties of ceria have been evaluated based on SOFC single cells running on methane fuel [6]. The single cell with ceria catalyst shows evidently improved performance in the initial operation for 14 h in methane fuel. Carbon deposition phenomenon, however, can still be identified in the following 6 h measurement. In order to enhance performance stability of single cells in hydrocarbon fuels,
80
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
research effort has been devoted toward optimization of the catalytic properties for ceria-based materials. Recent research results indicate that catalytic properties of ceria are essentially related to their crystal structures and defect properties [7e9]. Various elements, such as lanthanum, zirconium and calcium, have been doped into ceria to enhance the electrocatalytic property. Among them, the calcium-doped ceria shows highest electronic conductivity and oxygen ionic mobility, which are beneficial for electrochemical oxidation of hydrocarbon fuels. The doping of calcium into CeO2 produces lattice oxygen vacancy by a charge-compensating mechanism [10]:
ð1 xÞCeO2 þ xCaO/Ce1x Cax O2x þ zVO,,
(2)
The introduction of calcium increases the oxygen transport properties of ceria through the creation of the oxygen vacancy defects, favoring the conduction of oxygen ions. The increased oxygen ion mobility could facilitate the oxidation of hydrocarbon fuels, suppressing the carbon deposition in the anode. In addition, the doping of calcium will increase the lattice parameters of ceria, thus it has been used as a promoter for oxygen storage capacity of ceria. Compared to undoped ceria, the calcium-doped ceria demonstrates a higher oxygen storage capacity (~40% higher than CeO2), and it displays a remarkably increased redox activity [11,12]. Due to these promising results, it has been used as a catalyst for reforming of hydrocarbon fuels [13,14]. As calcium-doped ceria exhibits both good oxygen ionic conductivity and some catalytic property for hydrocarbon fuel reforming, applying the material in the anode of SOFC would make the best use of its electrochemical function and catalytic characteristic. In this case, different oxygen sources (e.g., oxygen ions from the cathode, and oxygen from steam) in the anode can be effectively utilized to oxidize carbon species and suppress the carbon deposition problem. On the other aspect, the oxygen vacancy can be trapped by the defect clusters at high levels of element doping, lowering the conductivity of oxygen ions [15]. Thus further research is necessary to optimize the material composition for improving the electrochemical property. In this work, Ce1-xCaxO2-d (0 x 0.30) powders are synthesized by a urea combustion method. The crystal structure and conductivities of Ce1-xCaxO2-d are investigated with respect to the amount of calcium content. Catalytic properties of Ce1-xCaxO2-d are evaluated based on Ni-Ce1-xCaxO2-d anode supported single cells running on methane fuel. The optimum calcium doping amount in Ce1-xCaxO2-d is determined with respect to conductivity of Ce1xCaxO2-d ceramics, electrochemical performance and performance stability of the single cells running on methane fuel. 2. Material and methods 2.1. Powder preparation The Ce1-xCaxO2-d (0 x 0.30) powders were synthesized by a urea combustion method. Reagent grade Ca(NO3)2$4H2O, Ce(NO3)3$6H2O and CO(NH2)2 (99.99% purity from Sigma-Aldrich) were used as starting chemicals. The nitrates were weighed according to the nominal compositions and dissolved into deionized water to form a solution. CO(NH2)2 was added into the solution as a combustion agent. The mole ratio of CO(NH2)2 to the total metal cation content was determined to be two [16]. The solution was stirred for 1 h and heated on a hot plate until auto-ignition and self-sustaining combustion occurred. The obtained ash was subsequently grounded for 24 h in ethanol medium using a ball milling process with zirconia balls. Then the powder was calcined at 900 C for 2 h in air. Ce0.8Sm0.2O1.9 (SDC) electrolyte powder and NiO-Ce1-xCaxO2d anode powder with 60 wt% NiO and 40 wt% Ce1-xCaxO2-d were
synthesized by the same urea combustion process. La0.6Sr0.4Co0.2Fe0.8O3-d (LSCF) cathode powder was synthesized by a glycinenitrate combustion method. The detailed synthetic procedures have been published previously [16,17]. 2.2. Single cell fabrication Anode supported single cell with a configuration of porous NiOCe1-xCaxO2-d anode supporter/yttira-stabilized zirconia (YSZ)/SDC bi-layer electrolyte/LSCF cathode was fabricated by a dry-pressing and spin-coating process. The NiO-Ce1-xCaxO2-d was mixed with a pore former (corn starch, Sigma-Aldrich) at 15 wt% by a ball-milling process. The mixed powder was compressed into disc pellets at a pressure of 80 MP using a hydraulic press machine. The diameter of green pellets was 21 mm and the thickness was ~2.5 mm. The pellets were calcined at 1000 C for 3 h in air to obtain sufficient mechanical strength for the following spin-coating processes. 17 wt% YSZ powder (Tosoh, Japan) was mixed with a lab-made organic carrier to form a slurry for the spin-coating process. The organic carrier was composed of B73210 organic binder (30 wt%, Ferro Electronics Materials), terpineol (40 wt%, Sigma-Aldrich) and ethanol (30 wt%, Decon Labs Inc). The slurry was ball-milled for 24 h and spin-coated on the surface of the NiO-Ce1-xCaxO2-d substrate at a rotation speed of 4500 rpm for 30 s. The spin-coating process was repeated for 3 times in order to achieve a ~10 mm YSZ layer. To avoid the formation of pin holes and microcracks in the electrolyte, the layer was calcined at 500 C in air for 30 min before the deposition of another layer. The SDC layer was fabricated on the surface of YSZ layer in the same way. The NiO-Ce1-xCaxO2d anode substrate and YSZ/SDC bi-layer electrolyte was co-sintered at 1400 C for 4 h in air. After the co-sintering process, the anode supported half cell showed a diameter of 17 mm and a thickness of ~2 mm. The LSCF cathode layer with an area of 0.79 cm2 was deposited on the surface of SDC electrolyte by the similar spin-coating process. 14.1 wt% LSCF powder was dispersed in an organic slurry consisting of 85.1 wt% ethanol and 0.8 wt% ethyl cellulose (SigmaAldrich) by the ball-milling process. The cathode slurry was spincoated on the surface of SDC electrolyte at a rotation speed of 3000 rpm. Then, the cathode layer was sintered at 1000 C for 2 h in air. 2.3. Structural characterization Crystal phases of the Ce1-xCaxO2-d powders and ceramics were studied by the X-ray diffractometry (XRD, Rigaku Smart-Lab XRD) using Cu Ka radiation at 40 kV and 44 mA. The data were collected in the continuous mode in the 2q range of 20e80 at 0.5 min1. Microstructures of the sintered Ce1-xCaxO2-d ceramics and the NiCe1-xCaxO2-d anode supported single cells were investigated by a scanning electron microscope (SEM, FEI Quanta 450). 2.4. Conductivity measurement Conductivities of Ce1-xCaxO2-d ceramics were measured by direct current (DC) four-probe technique and alternating current (AC) impedance measurement (Metrohm Autolab 302N Electrochemical workstation with FRA32 Impedance Analyzer), respectively. For the DC four-probe measurement, a dense rectangular bar sample (length: 25.0 mm, width: 5.4 mm, and height: 4.0 mm) was prepared and four parallel silver-line electrodes were used in the DC conductivity test. In the case of AC impedance measurement, a circular disc sample (diameter: 8.8 mm and thickness: 1.0 mm) was used with both sides painted with silver paste. The test was carried
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
out at the input sinuous voltage of 20 mV in the frequency range of 1 Hze10 MHz.
81
d powder is investigated by image analysis method using Image
Electrochemical performance of the single cell with the Ni-Ce1anode was characterized using a lab-designed fuel cell testing system. The system consists of an Autolab electrochemical workstation (Metrohm 302N and 20 A current booster), a set of mass flow controllers (P4B, MKS Instruments) and alumina sample holders. An alumina-based ceramic binder (AREMCO Products) was used as the sealing material. The single cell was heated up to the testing temperature at 3 C min1 and a hydrogen flow rate of 100 ml min1 was used to reduce the nickel oxide in the anode, while the cathode side was exposed to air. After the open-circuit voltage (OCV) of the single cell reached a stable state, electrochemical performance of the single cell was tested in hydrogen as fuel. Then, the anode fuel was switched from hydrogen to humidified methane. A methane flow rate of 100 ml min1 was fed into a water bubbling system at the room temperature (25 C) to make humidified methane. The humidified methane was delivered to the anode of the single cell to study the electrochemical performance. Electrochemical impedance spectra of the single cell were obtained under the OCV condition in the frequency range of 0.01 Hze100 kHz. The amplitude of the input sinuous signal was set to be 10 mV for the impedance measurement. Performance stability of the single cell was evaluated at 800 C in humidified methane fuel as well. The methane flow rate was set to be 100 ml min1 in the anode, and the cathode was exposed to air. The performance stability was tested by monitoring voltages of the single cell at a constant current load of 500 mA cm2.
pro-plus 6.0 software [19]. For each sample, three SEM images from different parts of the powder are used to estimate the particle size distribution and the average particle size is determined to be ~300 nm by Guass fitting. The Ce1-xCaxO2-d powders with different calcium contents show similar morphology and particle size distribution in this work. Fig. 3 (a) shows XRD data of Ce1-xCaxO2-d (0 x 0.15) ceramics sintered at 1400 C. The ceramics exhibit single cubic fluorite phase by the XRD analyses. Fig. 3 (b) shows the short range (28e34 ) XRD spectra data. With the increase of calcium doping amount, diffraction peaks corresponding to (111) and (200) planes of the cubic fluorite structure shift to lower 2q ranges. The Rietveld refinement technique has been applied to evaluate the crystallographic parameters of the cubic Ce1-xCaxO2-d ceramics. The lattice parameters are listed in Fig. 3 (c). The lattice parameters generally increase with calcium content in Ce1-xCaxO2-d. This phenomenon can be ascribed to the larger ion radius of Ca2þ (1.14 Å) compared with that of Ce4þ (1.01 Å) [9]. Moreover, the variation of the lattice parameter basically obeys the Vegard's rule in this work [20]. Fig. 4 shows SEM images of Ce1-xCaxO2-d (0 x 0.15) ceramics sintered at 1400 C in air. The samples exhibit generally dense microstructure with few isolated pores. Archimedes' method has been used to estimate the density of the ceramics and the relative densities are found to be over 95% for all the samples under this investigation [21]. Moreover, the relative density displays an increasing trend with the calcium amount (shown in Supplementary Material Fig. S1), which is in agreement with the reduced pores observed in the Ce1-xCaxO2-d ceramics. Thus, for the calcium-doped ceria compounds, the calcium acts as a sintering agent [22].
3. Results and discussion
3.2. Conductivity
3.1. Structure characterization
AC impedance measurement has been carried out using Ce1(x ¼ 0.00, 0.05, 0.10 and 0.15) ceramic disc samples sintered at 1400 C. Fig. 5 shows Nyquist plots of the Ce1-xCaxO2-d ceramics obtained at 300 C in air. The Nyquist plots are characterized by high frequency semi-circles and a low frequency line feature. To distinguish resistances from grains and grain boundaries of Ce1CaxO2-d ceramics, the impedance spectra are fitted by an equivax lent circuit model ðR1 Q1 ÞðR2 Q2 ÞðR3 Q3 Þ (shown in the insert of Fig. 5(a)). R is the resistance and Q represents the constant phase element. The impedance ZQ of a constant phase element Q and the equivalent capacitance C can be calculated according to the following equations [23,24]:
2.5. Cell performance measurement
xCaxO2-d
Fig. 1 shows XRD patterns of Ce1-xCaxO2-d powders calcined at 900 C for 2 h in air. The characteristic diffraction peaks for the fluorite phase are identified with the calcium doping amount in the range of 0e30 mol.% (0 x 0.30). At high calcium doping levels (x ¼ 0.20 and 0.30), an additional small peak is observed, which is attributed to the CaO phase in the composite [11,18]. The results indicate a solubility limit of 15 mol.% for calcium-doped ceria materials, which is in agreement with previous published results [11]. Fig. 2 (a) shows SEM image of the Ce0.9Ca0.1O2-d powder calcined at 900 C in air. The powder consists of fine particles with some agglomerations. Particle size distribution of the Ce0.9Ca0.1O2-
xCaxO2-d
ZQ ¼ 1
Q ði$2pf Þn
. C ¼ ðR$Q Þ1=n R
Fig. 1. XRD patterns of Ce1-xCaxO2-d (0 x 0.30) powders calcined at 900 C in air.
(3) (4)
where i is the imaginary unit, f is the frequency, and n is the an exponent parameter. Table 1 lists the fitting results obtained from AC impedance spectra. Generally, the equivalent capacitance C1 for ðR1 Q1 Þ is on the orders of magnitude of 1011 F, thus, R1 can be attributed to the resistance from grains of Ce1-xCaxO2-d ceramics. While, the values of C2 are on the orders of magnitude of 108 F. Therefore, R2 can be regarded to be the grain boundary resistance. The equivalent capacitances from ðR3 Q3 Þ component vary in the range of 106-104 F, the large capacitances are attributed to the electrode processes on both surfaces of Ce1-xCaxO2-d samples [22,25]. The sample without Ca doping shows one discernable semicircle, and the total resistance (sum of grain and grain boundary
82
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
Fig. 2. (a) SEM image of Ce0.9Ca0.1O2-d powder calcined at 900 C in air and (b) Particle size distribution of the Ce0.9Ca0.1O2-d powder.
Fig. 3. XRD patterns of Ce1-xCaxO2-d (0 x 0.15) ceramics sintered at 1400 C in air: (a) 2q range of 20e80 and (b) 2q range of 28e34 . (c) Lattice parameters obtained from XRD analyses.
resistances) is determined to be 45,129 U by impedance fitting (shown in Fig. 5 (a)). With the doping of calcium, grain and grain boundary resistances can be resolved from the impedance spectra (shown in the Nyquist plots in Fig. 5 and in the impedance fitting results in Table 1). For all the samples under the investigation, the gain resistance accounts as major contribution to the total resistance of the ceramics. Thus, the variation of the conductivity (or resistance) can be ascribed to the effects of calcium doping on the ceria-based materials. When increasing the calcium doping amount from 0 mol.% to 10 mol.%, the grain resistance displays a significant reduction. On the one hand, partially replacing tetravalent cerium with divalent calcium could give rise to the formation of oxygen vacancies: CeO2
00
CaO!CaCe þ VO,, þ OxO
(5)
The increased number of charge carriers improves the conductivities of the ceramics and reduces the grain resistances. On the
other hand, the introduction of calcium leads to the expansion of crystal structure (Fig. 3 (c)), enhancing mobility of the oxygen ions [13,22]. The contribution factors from these two aspects could lower resistances of the ceramics. Additionally, the grain boundary resistance exhibits a decreasing trend with calcium doping as well. This could be attributed to the increased relative density of the ceramics (shown in Supplementary Material Fig. S1). Calcium has been reported to be a sintering promoter for ceria-based materials [22]. The increment of relative density could facilitate the transport of charge carriers at grain boundaries, and enhance total conductivities of the ceramics. At a higher calcium doping amount (e.g., 15 mol.%), the grain resistance displays an increasing trend (from 6054 to 6453 U at 300 C). This phenomenon could be associated with the formation of defect clusters, resulting in the loss of movable oxygen vacancies in the ceramic [14,15]. As a result, despite the reduced resistance from grain boundaries, the total resistance of the composite
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
83
Fig. 4. SEM images of Ce1-xCaxO2-d ceramics sintered at 1400 C: (a) x ¼ 0.00, (b) x ¼ 0.05, (c) x ¼ 0.10 and (d) x ¼ 0.15.
Fig. 5. Nyquist plots of Ce1-xCaxO2-d ceramics at 300 C in air: (a) x ¼ 0.00, (b) x ¼ 0.05, (c) x ¼ 0.10 and (d) x ¼ 0.15. The equivalent circuit model is shown in the insert of (a).
Table 1 Resistances of Ce1-xCaxO2-d ceramics at 300 C obtained from impedance spectra fitting. Ce1-xCaxO2-d x x x x
¼ ¼ ¼ ¼
0.00 0.05 0.10 0.15
C1 (F) 1.64 2.45 1.87 1.91
1011 1011 1011 1011
R1 (U)
C2 (F)
R2 (U)
C3 (F)
45,129 12,054 6054 6453
e 3.73 108 2.96 108 2.76 108
e 2406 595 392
6.32 1.32 4.74 4.28
(Ce0.85Ca0.15O2-d) shows an increasing trend. Fig. 6 shows conductivities of calcium-doped ceria ceramics obtained by the four-probe method at 800 C in air. The data derived from AC impedance spectroscopy tests are also included in the figure. The undoped ceria presents fairly low conductivity (0.002 S cm1 at 800 C), which is in agreement with previous reported data on ceria materials [26]. With the introduction of calcium, the conductivity improves significantly and the 10 mol.%
R3 (U) 106 104 104 104
642,278 85,542 37,928 40,694
calcium-doped sample (Ce0.9Ca0.1O2-d) exhibits a highest conductivity (0.102 S cm1 at 800 C) among all the specimens investigated in this work. The conductivity is competitive compared with the widely used samarium-doped ceria (0.082 S cm1 at 800 C) and gadolinium-doped ceria (0.053 S cm1 at 750 C) materials [27,28]. The high conductivity is believed to be favorable for electrochemical oxidation of hydrocarbon fuels in the anode. In addition, calcium is a much cheaper and widely available material in
84
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
Fig. 6. Conductivity of Ce1-xCaxO2-d ceramics at 800 C in air.
Fig. 7. SEM image of Ni-Ce0.9Ca0.1O2-d anode supported single cell.
comparison with samarium and gadolinium. Application of the material could help reduce the cost of SOFCs. Thus, 5e15 mol.% calcium-doped ceria are applied into the anode of the single cell consisting of Ni-Ce1-xCaxO2-d anode supporter/YSZ/SDC bi-layer electrolyte/LSCF cathode. 3.3. Single cell performance Fig. 7 shows a typical cross-sectional SEM image of the single cell. Image analysis method is used to estimate thicknesses of cell functional layers. The LSCF cathode layer exhibits good adhesion to the SDC layer and the thickness is determined to be ~10 mm. The SDC buffer layer appears to be dense with some isolated pores. The YSZ electrolyte layer displays a dense microstructure and the thickness is ~9 mm. The Ni-Ce0.9Ca0.1O2-d anode substrate presents
highly porous microstructure, which could be due to the burning out of pore former in the cell fabrication process. Fig. 8 (a) shows voltage and power density of the single cell with Ni-Ce1-xCaxO2-d (x ¼ 0.05, 0.10 and 0.15) anodes at 800 C in humidified methane fuel. The single cells present similar OCV of ~1.02 V, indicating a dense electrolyte layer and good sealing for the single cell. The maximum power density exhibits a slight variation as a function of calcium content in the Ce1-xCaxO2-d. The NiCe0.9Ca0.1O2-d (x ¼ 0.10) anode supported single cell displays an optimum maximum power density of 618 mW cm2 at 800 C in methane fuel. The performance is comparable with the typical NiCe0.8Sm0.2O1.9 anode supported single cell (572 mW cm2 at 800 C in methane fuel, shown in Supplementary Material Fig. S2) under the same operation condition. Electrochemical impedance spectra of the single cells are obtained under the OCV condition in methane fuel. Fig. 8 (b) shows Nyquist plots of the single cell at 800 C. The impedance spectra are fitted by the LRohm ðRH QH ÞðRL QL Þ equivalent circuit model (shown in the insert of Fig. 8(b)). The inductance element L represents the inductive impedance response, while Rohm is the ohmic resistance of the entire cell. ðRH QH ÞðRL QL Þ components denote the electrochemical processes from both the anode and the cathode. The fitting results of the electrochemical impedance spectra are listed in Table 2. The equivalent capacitance for the higher frequency impedance arc is on the order of magnitude of 103-102 F cm2. The electrode processes can be ascribed to the charge transfer process ðRH QH Þ. The lower frequency arc has a larger equivalent capacitance (~100 F cm2), which can be attributed to the gas adsorption and dissociation, and gas phase diffusion processes ðRL QL Þ. The total polarization resistance of the whole cell ðRp Þ is the sum of RH and RL [29e31]. The single cells illustrate similar ohmic resistances (0.26e0.27 U cm2, at 800 C in Table 2) with different calcium contents in Ce1-xCaxO2-d. For the nickel containing anode supported single cell with YSZ as electrolyte, the ohmic resistance from the electrolyte can be regarded as a major contribution to the ohmic loss of the whole cell. The similar ohmic resistance indicates a comparable electrolyte layers for different cells in this research. As the fabrication process of the LSCF cathode layer is kept same in this work, the variation of the polarization resistance (Rp ) could be attributed to differences in the anode. By comparing variations of RH and RL in Table 2, the differences of the total polarization resistance (Rp ) are due to the charge transfer process (RH ). As the Ce0.9Ca0.1O2-d demonstrates a higher conductivity among all the samples investigated in this work (shown in Figs. 5 and 6), the good conductivity could enhance the transfer of oxygen between the anode and the fuel, thus, improving the charge transfer processes at the anode [13]. This data is in agreement with the higher maximum power density obtained for the cell using Ni-Ce0.9Ca0.1O2-d as the anode.
Fig. 8. (a) Cell voltage and power density of Ni-Ce1-xCaxO2-d (x ¼ 0.05, 0.10 and 0.15) anode supported single cells at 800 C in methane, (b) Nyquist plots of the single cells under OCV condition, and (c) the performance stability. The equivalent circuit model is shown in the insert of (b).
K. Zhao, Y. Du / Journal of Power Sources 347 (2017) 79e85
85
Table 2 Electrochemical parameters of the single cells at 800 C in methane. Ni-Ce1-xCaxO2-d anode
Rohm (U cm2)
CH (F cm2)
RH (U cm2)
CL (F cm2)
RL (U cm2)
Rp (U cm2)
x ¼ 0.05 x ¼ 0.10 x ¼ 0.15
0.27 0.26 0.27
1.51 102 9.83 103 1.29 102
0.20 0.12 0.17
1.81 2.12 1.65
0.12 0.10 0.11
0.32 0.22 0.28
Fig. 8 (c) shows performance stability of the Ni-Ce1-xCaxO2-
Appendix A. Supplementary data
d (x ¼ 0.05, 0.10 and 0.15) anode supported single cells in methane
fuel at 800 C. The performance stability is investigated under the galvanostatic mode (at the current density of 500 mA cm2). Among all the single cells investigated in this research, the NiCe0.9Ca0.1O2-d anode supported cell displays a lowest degradation rate (1.1 mV h1) during 30 h operation in methane fuel. The degradation rate is comparable with typical Ni-Ce0.8Sm0.2O1.9 anode supported single cell (~1.6 mV h1) in the initial 10 h operation (shown in Supplementary Material Fig. S2). However, the overall stability for the Ni-Ce0.9Ca0.1O2-d anode supported cell is evidently enhanced compared with the Ni-Ce0.8Sm0.2O1.9 supported cell. After the performance stability test, SEM/EDX has been used to estimate the amount of carbon in the anode. Only a small amount of carbon (4.8 wt%) is detected in the Ni-Ce0.9Ca0.1O2d anode, while 16 wt% carbon is found out in Ni-Ce0.8Sm0.2O1.9 anode (shown in Supplementary Material Fig. S3 and Table S1). The higher tolerance of Ni-Ce0.9Ca0.1O2-d anode to carbon deposition suggests the application potential of calcium-doped ceria material (Ce0.9Ca0.1O2-d) as an anode component for SOFCs running on methane fuel. 4. Conclusions Calcium-doped ceria with nominal compositions of Ce1-xCaxO2d (0.00 x 0.30) is synthesized by the urea-combustion method. Crystal phases of the Ce1-xCaxO2-d powders are investigated with
respect to the calcium doping amount. The powders show single fluorite phase when the doping level is within 15 mol.% (x 0.15), and higher calcium doping leads to the formation of secondary CaO phase. Conductivities of the Ce1-xCaxO2-d ceramics are studied as a function of calcium amount as well. The 10 mol.% calcium-doped sample (Ce0.9Ca0.1O2-d) exhibits a high conductivity of 0.102 S cm1 at 800 C, which is competitive as compared with typical samarium- or gadolinium-doped ceria. The calcium-doped ceria is applied in the Ni-Ce1-xCaxO2-d anode supported single cell, and the cell with the Ni-Ce0.9Ca0.1O2-d anode gives an optimum maximum powder density (618 mW cm2 at 800 C) and reasonable performance stability during the ~30 h test in methane fuel. The promising electrochemical results suggest the application feasibility of the Ce0.9Ca0.1O2-d materials as the anode component of SOFCs. Acknowledgments The authors acknowledge the support from Kent State University and NASA (Contract No.: NNX14CC37P and NNX15CC12C).
Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jpowsour.2017.01.113. References [1] D. Sarantaridis, A. Atkinson, Fuel Cells 7 (2007) 246e258. [2] M.S. Khan, S.B. Lee, R.H. Song, J.W. Lee, T.H. Lim, S.J. Park, Ceram. Int. 42 (2016) 35e48. [3] N. Mahato, A. Banerjee, A. Gupta, S. Omar, K. Balani, Prog. Mater. Sci. 72 (2015) 141e337. [4] A. Trovarelli, Cat. Rev. 38 (1996) 439e520. [5] M. Ozawa, M. Kimura, A. Isogai, J. Alloys Compd. 193 (1993) 73e75. [6] Y. Yin, S. Li, C. Xia, G. Meng, J. Power Sources 167 (2007) 90e93. [7] N. Jaiswal, S. Upadhyay, D. Kumar, O. Parkash, J. Power Sources 222 (2013) 230e236. [8] C. Zhang, X.D. Wen, B.T. Teng, Y. Zhao, M. Fan, Fuel Process. Technol. 131 (2015) 1e6. € fer, Acta Mater. 103 (2016) 361e369. [9] A. Maheshwari, H.D. Wiemho [10] Y. Zheng, Y. Shi, H. Gu, L. Gao, H. Chen, L. Guo, Mater. Res. Bull. 44 (2009) 1717e1721. [11] V. Thangadurai, P. Kopp, J. Power Sources 168 (2007) 178e183. [12] Y. Zhang, S. Andersson, M. Muhammed, Appl. Catal. B Environ. 6 (1995) 325e337. [13] H. Song, U.S. Ozkan, J. Phys. Chem. A 114 (2010) 3796e3801. jcik, [14] S. de Carolis, J.L. Pascual, L.G.M. Pettersson, M. Baudin, M. Wo K. Hermansson, A.E.C. Palmqvist, M. Muhammed, J. Phys. Chem. B 103 (1999) 7627e7636. [15] M. Mogensen, T. Lindegaard, U.R. Hansen, G. Mogensen, J. Electrochem. Soc. 141 (1994) 2122e2128. [16] Q. Xu, D.P. Huang, K. Zhao, W. Chen, M. Chen, B.H. Kim, Ceram. Int. 37 (2011) 913e920. [17] K. Zhao, Q. Xu, D.P. Huang, M. Chen, B.H. Kim, Ionics 17 (2011) 247e254. [18] M. Yan, T. Mori, J. Zou, F. Ye, D.R. Ou, J. Drennan, Acta Mater. 57 (2009) 722e731. [19] M. Rochoux, Y. Guo, Y. Schuurman, D. Farrusseng, Phys. Chem. Chem. Phys. 17 (2015) 1469e1481. [20] S. Zha, C. Xia, G. Meng, J. Power Sources 115 (2003) 44e48. [21] K.Y. Park, T.H. Lee, J.T. Kim, N. Lee, Y. Seo, S.J. Song, J.Y. Park, J. Alloys Compd. 585 (2014) 103e110. [22] S. Banerjee, P.S. Devi, D. Topwal, S. Mandal, K. Menon, Adv. Funct. Mater. 17 (2007) 2847e2854. [23] C. Zhang, S. Wang, R. Guo, L. Liu, Y. Yang, J. Alloys Compd. 646 (2015) 283e290. [24] M. Zunic, G. Brankovic, C.R. Foschini, M. Cilense, E. Longo, J.A. Varela, J. Alloys Compd. 563 (2013) 254e260. [25] K.C. Anjaneya, G.P. Nayaka, J. Manjanna, G. Govindaraj, K.N. Ganesha, J. Alloys Compd. 578 (2013) 53e59. [26] S.R. Bishop, T.S. Stefanik, H.L. Tuller, J. Mater. Res. 27 (2012) 2009e2016. [27] M. Chen, B.H. Kim, Q. Xu, B.K. Ahn, W.J. Kang, D.P. Huang, Ceram. Int. 35 (2009) 1335e1343. [28] A.Z. Liu, J.X. Wang, C.R. He, H. Miao, Y. Zhang, W.G. Wang, Ceram. Int. 39 (2013) 6229e6235. [29] C. Lalanne, F. Mauvy, E. Siebert, M.L. Fontaine, J.M. Bassat, F. Ansart, P. Stevens, J.C. Grenier, J. Eur. Ceram. Soc. 27 (2007) 4195e4198. [30] E.C. Shin, P.A. Ahn, H.H. Seo, J.M. Jo, S.D. Kim, S.K. Woo, J.H. Yu, J. Mizusaki, J.S. Lee, Solid State Ionics 232 (2013) 80e96. [31] V.C. Kournoutis, F. Tietz, S. Bebelis, Fuel Cells 9 (2009) 852e860.