Scholarly article on topic 'Enhanced electrochemical activity in Ca3Co2O6 cathode for solid-oxide fuel cells by Cu substitution'

Enhanced electrochemical activity in Ca3Co2O6 cathode for solid-oxide fuel cells by Cu substitution Academic research paper on "Materials engineering"

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{"Solid-oxide fuel cells" / "Ca3Co2O6 cathode" / "Cu substitution" / "Electrochemical performance"}

Abstract of research paper on Materials engineering, author of scientific article — Fushao Li, Rui Zeng, Long Jiang, Tao Wei, Xuefei Lin, et al.

Abstract Cu-substituted Ca3(Co1−x Cu x )2O6 (x = 0, 0.05, 0.1, 0.15) are prepared and evaluated as cathode materials for solid-oxide fuel cells (SOFCs). Effects of Cu substitution for Co on structure, electrical conductivity, thermal expansion and electrochemical performance have been investigated. Pure hexagonal structure can be attained, but a small amount of impurity phase Ca0.828CuO2 appears if x is high. Ca3Co2O6 behaves as a thermal-activated semiconductor in the temperature range between 300 and 800 °C; its electrical conductivity is remarkably enhanced by Cu substitution. Experimental results show coexistence of mixed valent Co ions and disordered state of oxygen vacancies in the Cu-substituted samples. Meanwhile, Cu substitution leads to slightly enlarged thermal expansion coefficient, reduced area specific resistance and improved electrochemical performance. For x = 0.05, the power density as high as ca. 550 mW cm−2 at 800 °C is achieved in the single cell with La0.8Sr0.2Ga0.83Mg0.17O2.815 as electrolyte and Ni–Ce0.8Sm0.2O1.9 as anode. Cu substitution can effectively enhance the electronic and ionic transport properties of Ca3Co2O6 cathode for SOFCs.

Academic research paper on topic "Enhanced electrochemical activity in Ca3Co2O6 cathode for solid-oxide fuel cells by Cu substitution"


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Journal of Materiomics 1 (2015) 60—67

Enhanced electrochemical activity in Ca3Co2O6 cathode for solid-oxide

fuel cells by Cu substitution

Fushao Lia,b, Rui Zeng b, Long Jiang b, Tao Weib, Xuefei Lin a, Yingxian Xu a, Yunhui Huang b *

a School of Chemistry and Chemical Engineering, Qujing Normal University, Qujing, Yunnan 655011, China b School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan, Hubei 430074, China

Received 5 January 2015; revised 19 January 2015; accepted 30 January 2015 Available online 4 April 2015


Cu-substituted Ca3(Co1_xCux)2O6 (x — 0, 0.05, 0.1, 0.15) are prepared and evaluated as cathode materials for solid-oxide fuel cells (SOFCs). Effects of Cu substitution for Co on structure, electrical conductivity, thermal expansion and electrochemical performance have been investigated. Pure hexagonal structure can be attained, but a small amount of impurity phase Ca0 828CuO2 appears if x is high. Ca3Co2O6 behaves as a thermal-activated semiconductor in the temperature range between 300 and 800 °C; its electrical conductivity is remarkably enhanced by Cu substitution. Experimental results show coexistence of mixed valent Co ions and disordered state of oxygen vacancies in the Cu-substituted samples. Meanwhile, Cu substitution leads to slightly enlarged thermal expansion coefficient, reduced area specific resistance and improved electrochemical performance. For x — 0.05, the power density as high as ca. 550 mW cm~2 at 800 °C is achieved in the single cell with La08Sr0.2Ga0.83Mg0.17O2815 as electrolyte and Ni—Ce0.8Sm0.2O1.9 as anode. Cu substitution can effectively enhance the electronic and ionic transport properties of Ca3Co2O6 cathode for SOFCs.

© 2015 The Authors. Production and hosting by Elsevier B.V. on behalf of The Chinese Ceramic Society. This is an open access article under the CC BY-NC-ND license (

Keywords: Solid-oxide fuel cells; Ca3Co2O6 cathode; Cu substitution; Electrochemical performance

1. Introduction

Transition-metal oxides are commonly used as catalysts for oxygen reduction reaction (ORR) in solid-oxide fuel cells (SOFCs) due to their excellent chemical stability and oxidation resistance at elevated temperatures [1 — 3]. Among them, cobalt-containing ABO3 perovskite compounds, such as Baj_xSrx-


1 -xSr^Co! _yFeyO3 _8, Sr: _xRExCoO3 _8 (RE = rare earth) and their derivatives, have been intensively investigated in the last decades since they generally show prominent electro-catalytic activity and impressive electrical conductivity at the reduced operating temperatures (500 -800 °C) [4—11]. However, most of the cobaltite-based perov-skites exhibit rather large thermal expansion coefficient (TEC),

* Corresponding author. E-mail address: (Y. Huang). Peer review under responsibility of The Chinese Ceramic Society.

which renders them to have poor "weldability" with those popular oxygen-ion conductive electrolytes like Y2xZrj_2xO2_x (YSZ), Ce1_xSmxO2_8 (SDC) and La1_xSrxGa1_yMgyO3_8 (LSGM) during high-temperature sintering [12]. An immediate solution to this problem is to tune the sublattice composition with substituting ions. But in most cases, this effect is almost negligible compared with the big difference in TEC between them. On the other hand, once the dopant is introduced, discrepancy in ion radius between the host and guest atoms can only be counteracted either by tilting, displacement or cooperative rotation of the M-O polyhedron subcells, and consequently structural distortion is almost inevitable. In some extreme cases, the first-order phase transition or collapse of the original lattice may then happen, which will dramatically decrease the electronic and ionic transport properties. Furthermore to some doped perovskite compounds, segregation or enrichment of surface elements is another uncontrollable problem, such as Sr-enrichment on Laj_xSrxCoO3_5 surface [13,14]. Therefore, it is natural to think about that, if those key

2352-8478/© 2015 The Authors. Production and hosting by Elsevier B.V. on behalf of The Chinese Ceramic Society. This is an open access article under the CC BY-NC-ND license (

electrochemical properties of Co-based perovskites can still be retained, finding other Co-containing structures with likewise stable lattices would be another alternative pathway.

In this regard, Ca3Co2O6 (CCO) is a good candidate that was recently explored as a promising cathode material in our group [15]. As a member of large fascinating family A3„+3mA'nB3m+nO9m+6„ [16], Ca3Co2O6 (m = 0, n = 1, A' = B) crystallizes into hexagonal closed-packing arrays (R-3c space group) with pseudo one dimensional feature. The Ca atoms are evenly distributed around [Co2O6]f pillars which consist of alternative face-shared CoO6 octahedra and distorted CoO6 trigonal prisms (Fig. 1). However, the electrical conductivity of CCO is low (5-9 S cm-1 at 500-800 °C), which is unfavorable for the electrochemical performance. In the present study, to address this problem, Cu2+ was chosen as the dopant in CCO to improve the transport properties and hence the electrochemical performance.

Since Cu usually exists as a stable bivalent state, a low level substitution at the Co-sites is expected to improve the transport properties of CCO by enhancing the carrier concentration [17—19]. With Cu2+ incorporation into Co-sublattice sites, to retain electro-neutrality the excessive negative charges must be compensated either by creation of oxygen vacancies or by generation of holes, i.e. oxidation of Co3+ to Co4+, as illustrated respectively in Kroger—Vink note equations (1) and (2):

Ca3Co2O6 ,

2CuO -3—2o6 2CuCo + 2OX + VO

^ Ca3Co2O6 - v / \

CuO + Co^0 -- Cuc0 + CoC0 + O0 (2)

Therefore, Cu substitution is hopeful to enhance the electrochemical properties of CCO cathode.

Accordingly, Cu-substituted CCO compounds were synthesized via a sol—gel route and evaluated as cathode materials for SOFCs. Low-level Cu substitution in Ca3(Co1-xCux)2O6 was carried out from x = 0 to 0.15 to optimize the electrochemical performance.

2. Experimental

Polycrystalline Ca3(Co1-xCux)2O6 powders were prepared

with citrate complex as precursor. As starting reagents, stoi-

chiometric Ca(NO3)2 • 4H2O, C4H6O4Co$4H2O (cobalt

acetate) and Cu(NO3)2 • 3H2O in analytical grade were dissolved into a small amount of de-ionized water under continuous stirring. Citric acid and ethylene glycol were then added as chelating agent and disperser, respectively. The molar ratio of total metal ions to citric acid was 1:2. The mixture of solution was then put on a plate heater to evaporate till a homogeneous gel was obtained. The obtained gel was first decomposed at 400 °C in air for 5 h, and then annealed at 900 °C in air for 10 h after grinding. The resultant powders were further treated according to the requirements of characterization items. La0.8Sr02Ga0.83Mg0.17O2.8i5 (LSGM) dense electrolyte discs and Ce08Sm02Oi.9 (SDC) powders were prepared via solid-state reaction, as described in our previous works [20,21]. The anode composite was made by thoroughly mixing NiO and SDC fine powders in weight ratio of 65:35.

The phase of the samples was detected by X-ray powder diffraction (XRD, Philips X'Pert PRO diffractometer) in Bragg-Brentano reflection geometry with Cu Ka radiation at 40 kV and a receiving slit of 0.2—0.4 mm. The diffraction patterns were collected at room temperature (RT) by step scanning in the range of 10° < 2d < 80° with a scan rate of 2° min-1. The binding state of compositional elements in Ca3(Co1-xCux)2O6 was analyzed by using a MULT1LAB2000 X-ray Photoelectron Spectrometer. The incident radiation was monochromatic Al Ka X-rays (1486.6 eV). Narrow highresolution scans were run to obtain O1s and Co2p level spectra with 0.05 eV steps. All binding energies were referenced to the C1s peak (285 eV) arising from adventitious carbon.

The oxygen nonstoichiometry was measured by thermog-ravimetric analysis (TGA) (Stanton STA 781 instrument) carried out under ambient pressure in a temperature range from RT to 1000 °C. The amount of sample powder was ~15 mg and the heating rate was 10 °C min-1. The thermal expansion was measured on the rectangular-shaped bar samples (5 mm o 5 mm o 20 mm) from RT to 1000 °C at a heating rate of 10 °C min-1 by using a dilatometer (NETZSCHSTA449c/3/G). Temperature dependence of electrical conductivity was studied with a modified four-probe method on the manually sampled RTS-8 digital instrument in the stagnant air. To prepare specimens for the conductivity measurement, the as-prepared sample powders were pressed into a pellet with diameter of 13 mm and thickness of 1 mm

Fig. 1. Crystal structure of Ca3Co2O6 (left: hexagon cell; right: projected view along c-axis).

under a pressure of 100 MPa followed by sintering at 900 °C for 10 h.

Single cells were fabricated with an electrolyte-supported technique. LSGM disc with thickness of 230(±5) mm was used as electrolyte, NiO-SDC composite as anode, and SDC as the buffer layer between the anode and the electrolyte to prevent inter-diffusion of cationic species. The SDC slurry was first screen-printed onto one side of LSGM disc and baked at 1300 °C for 2 h in stagnant air, followed by screen-printing the "NiO-SDC'' slurry (~20 mm) and successively baking at 1250 °C for another 4 h. The cathode slurry was screen-printed onto another side of the electrolyte and fired at 950 °C for 5 h to achieve porous texture with layer thickness of ~20 mm. Ag grids and Ag paste were used as current collectors to minimize the catalytic effect. The cell was mounted onto an alumina tube and meticulously sealed by Ag paste. H2 was fed to the anode side as a fuel at a flow rate of 60 ml min_: while the cathode side was exposed to the ambient air. Electrochemical tests were performed on a 2273 electrochemical workstation driven by software package PowerSuite.

Symmetric cells with conformation of "cathode | LSGM | cathode" were employed to test electrochemical impedance spectroscopy (EIS). The fabrication method of the symmetric cells is almost same as that of single cells except that the cathode slurry was simultaneously screen-printed onto the both sides of LSGM disc. EIS across symmetric cells were collected around open current voltage (Eocv) using a voltage amplitude of 10 mV and a frequency range from 100 kHz to 10 mHz. The program of ZsimpWin was used to construct the equivalent circuit models and to simulate the EIS measurement based on the least-squares method. I—V curves were measured on the single cells to demonstrate the power density output.

3. Results and discussion

The phase of the synthesized samples was checked by XRD. Phase-pure Ca3Co2O6 with hexagonal crystal structure can be easily attained by annealing at 850_950 °C. The reaction temperature has little effect on the solubility of Cu2+ ions in CCO. The XRD patterns at RT are presented in Fig. 2 for Ca3(Co:_xCux)2O6 (x = 0, 0.05, 0.1, 0.15) samples obtained by being annealed twice at 900 °C for 10 h with

a | nag □ Caos21Cu02-PDF#48-0212 II . 11 .....i .

1 1 *=0.10

1 1 II II ... .,

1. 1 1. .11 .. .,. *T°.00

1. . 1 Ca Co O -PDF#51-0311 3 2 6


JF-0.00 jy

thorough intermittent grinding. All the patterns are indexed to the diffractions of the hexagonal lattices. However, a minority of impurity phase, discerned as Ca0.828CuO2, begins to appear in the sample with x = 0.1, and becomes more detectable at x = 0.15. For simplicity, all the compounds are still referred as Ca3(Co1_xCux)2O6 despite the existence of secondary phase and intrinsic non-stoichiometry.

Back to Fig. 2, the XRD patterns also indicate that, once x = 0.05 mol fraction of Cu substitution are introduced, the main diffraction peaks shift toward lower angles compared with the undoped one (as shown in Fig. 2b), indicative of successful incorporation of Cu-ions into the lattice of Ca3Co2O6. To be more quantitative, the lattice parameters calculated by the Jade software are listed in Table 1. Upon Cu substitution, both lattice parameter a and c increase, especially along c-direction. Since Cu2+ (0.73 A) has more stable valence and larger ionic radius relative to Co3+ (0.55 A for low spin and 0.61 A for high spin), the expanded lattice is ascribed to two factors: one is the substitution of Cu for Co, and the other is the cationic valence change due to the substitution. In view of ion transport, the expanded lattice volume by Cu2+ is expected to support higher degree of freedom for oxygen ionic movement and in turn to facilitate the cathodic process of oxygen reduction reaction. Nevertheless, the peaks shift back from lower angles when the nominal Cu substitution level increases to x = 0.1 and 0.15. We surmise that this phenomenon may result from A-site deficiency when a small amount of non-stoichiometric Ca2+ is consumed by the impurity phase. For the cathode materials, the A-site deficiency is important to the electrochemical performance.

Chemical states of O and Co are crucial to determine the ionic and electronic transport properties of the present oxides. The XPS curves of O 1s and Co 2p are shown in Fig. 3. The O 1s spectrum of the undoped sample consists of two components at 528.5 and 531.2 eV (Fig. 3a). One should reasonably be assigned to the surface oxygen species (Osurf) while the other comes from the lattice ones (OLatt). Superimposed upon this profile, evolution of O 1s spectra with increasing Cu2+ content shows the obvious change in chemical environment of O due to entry of the guest ions. The x = 0.05 sample shows similar XPS profile to the undoped one. For x = 0.10 and 0.15, both position and intensity of the peaks change, demonstrating that some impurity phases appear.

Fig. 3b shows the asymmetric Co 2p core-level spectra of the Ca3(Co1_xCux)2O6 samples. For the undoped sample, two asymmetric peaks of Co 2p1/2 and Co 2p3/2 are obviously shaped; with Cu substitution, the peaks become wider. This indicates that the chemical environment changes due to Cu incorporation. By carefully checking, we can see some weak satellite peaks between the two main peaks, indicating that

Table 1

Hexagonal lattice parameters of Ca3(Co1_xCux)2O6 obtained from Jade program.

33 34 35

Fig. 2. XRD patterns of Ca3(Co1_xCux)2O6 synthesized at 900 °C: (a) survey inspection, (b) localized zoomed view.

Sample Space group a = b (A) c (A) a = b (°) g (°)

x = 0.00 R-3c 9.078 (2) 10.380 (3) 90 120

x = 0.05 R-3c 9.083 (2) 10.406 (4) 90 120

*=0.15 /

x=0.10 f Shoulder peak ca &

x=0.05 S *=0.00 a 1

525 530 535 Binding energy (eV)

0.05j \ ^^^ ML


770 775 780 785 790 795 800 805 Binding energy (eV)

Fig. 3. XPS of Ca3(Co1-xCux)2O6 at RT: (a) O1s; (b) Co 2p.

small amount of other valent Co ions may exist. These results demonstrate that Cu2+ substitution can effectively create electron holes in Ca3(Co1-xCux)2O6, which increases the carrier concentration.

The electrical conductivity as a function of temperature for Ca3(Co1-xCux)2O6 is shown in Fig. 4a. A thermally activated behavior is observed in all samples over the whole measured temperature range, exhibiting a typical semiconductor-like property. The x = 0.05 sample exhibits the maximum conductivity, which is more than two times higher than that of the un-doped one. With further increasing x, the conductivity begins to drop. Combined with XRD examination, it can be deduced that the incorporation of Cu2+ ions in the lattice can effectively enhance the electric transport, but the appearance of Ca0.828CuO2 impurity is disadvantageous to the conductivity.

In general, the cationic substitution has some effect on the thermal-activated conduction, but it is difficult to know how the substitution enhances the electrical conductivity [17]. In the present case, the conduction process may be through weakly localized or partly itinerant electronic defects. Assuming Ca3(Co1-xCux)2O6 oxides exhibit polaronic conduction that is governed by a small-polaron hopping mechanism with the expression:

a = —exp — T \kT

then a linear expression can be obtained by plotting log(sT) vs 1000/T. The plots are displayed in Fig. 4b. For the undoped sample, the plots are hardly linearly fitted over the

measured temperatures, meaning that the conductive nature in the low temperatures differs from that in the high temperatures. For the Cu-substituted samples, the plots fit well in a line, demonstrating that the conduction is dominated by some fixed mechanism. The difference may be associated with the highly anisotropic crystallographic feature of the structure. For the undoped sample, the conduction in the low temperatures takes place as the result of disproportional reaction, as shown in equation (4) [23]:

2Co0o / COco + CoCo

in this regard, conducting process is most likely confined completely within the separated parallel one-dimensional [Co2O6+]TC pillars along the hexagonal c-axis. But at high temperatures, the oxygen vacancies are generated, so the conducting path can be created in another mode as shown in equation (5) [23]:

O0+ Co0o/CoCo+2O2(g) + vO" (5)

In this case, inter-chain charge exchange is reinforced, and a new conduction mechanism then gets involved in. Based on these viewpoints, when the Cu-substituted sample is concerned, it is easy to imagine that the inter-chain charge exchange behavior prevails almost over the whole temperature range because, the oxygen vacancies have already been yielded even before heating due to incorporation of aliovalent Cu2+ (equation (2)). For this reason, these two conduction mechanisms simultaneously exist over the whole temperature range.

400 600

Temperature (°C)

"a 30 s

i2-5 "So

■s 2.0

0.8 1.0 1.2 1.4 1.6


Fig. 4. (a) Temperature dependent electrical conductivity (s), and (b) Arrhenius plots of log(sT) vs 1000/T for Ca3(Coi-xCux)2O6 in the air.

As a cathode material for SOFC, the property with mixed ionic and electronic conductivity is usually favored. In order to produce the mixed conductivity, there should be oxygen vacancies existing in the lattice and hence the ionic conduction may result from the hops among the vacancies. More oxygen vacancies can support higher transport rate of the oxygen ions and larger triple-phases boundary (TPB) active region for oxygen reduction reaction (ORR). If the doping level is extremely high, the oxygen vacancies may order as the constituting elements in a new structure, and they can then no longer be considered as defects and are not mobile by themselves [24]. TG measurements were performed to check whether the disorder-order transition occurs in the samples. As shown in Fig. 5, the undoped sample undergoes a progressive mass loss after a short initial gain, which may be caused by equilibrating process of the measuring apparatus. No obvious plateau is observed during the heating, suggesting that oxygen vacancies have no long-range order. The presence of disordered oxygen vacancies could account, by itself, for the necessary ionic conductivity of oxide anions that is required for the mass transport across the cathode [25]. For the x = 0.05 and x = 0.1 samples, only a very small mass lose occurs, and there is a seeming plateau on the curves. Based on the XPS analysis, this type of plateaus should not be ascribed to the ordering of oxygen vacancies but to the structural compactness of hexagonal closed-packing nature that allows only a very limited amount of oxygen anions to escape. Therefore, in our opinion, once a certain amount of oxygen vacancies are generated, it is difficult for the additional lattice oxygen ions to escape from the hexagonal closed-packing structure. As for the x = 0.15 sample, the mass loss looks too abnormal to be explained due to the relatively high content of impurity phase.

Thermal stability is closely associated with the elevated temperature during the operation of SOFCs. We measured thermal expansion of all samples in air from 25 to 1000 °C. As shown in Fig. 6, all curves indicate a pseudo linear expansion between 200 °C and 950 °C. According to calculation on differential linear expansion (AL/L0) versus temperature, the average TEC values within this temperature range for all the measured samples are 16-22 (10~6 K-1), which are

comparable with those of the commonly-used electrolytes. The TEC increases slightly with x (as the arrow marked in Fig. 6). Importantly, the continuous progressive change in thermal expansion also supports the above discussion that there is no oxygen ordering phase formed during the heating in all samples [22].

Kinetics of oxygen reduction reaction in cathode was studied by EIS across the electrolyte-supported symmetric cathode. The corresponding Nyquist plots of the cells are collected in Fig. 7a. All plots contain essentially two separable capacitance arcs arising from two relaxation processes of the electrode, which means that ORR is governed by at least two different electrode processes. Accordingly, the cathode-electrolyte systems can be analyzed with the equivalent circuit of Rohm(RiQi)(R2Q2), in which the overall ohmic resistance (Rohm) includes the electrolyte resistance, the electrode ohmic resistance and the lead resistance. Across the arcs, the high-frequency resistance is probably associated with chargetransfer process (R1); the low-frequency arc is ascribed to diffusion process (R2), including adsorption-desorption of oxygen, oxygen diffusion at the gas-cathode interface, and the surface diffusion of intermediate oxygen species [26,27]. The difference between the real axes intercepts of the impedance plot is considered as the cathode polarization resistance (Rp = R1 + R2), which is commonly treated as the area specific resistance (ASR). On the whole, the ASR at different temperature for the x = 0.05 sample is evidently reduced to some extent.

In addition, some other interesting phenomena can be observed from EIS data. Focusing on the Nyquist plots at 800 °C (Fig. 7b), we can see that the shapes of capacitive arcs are quite different from each another. For the undoped sample, it consists of a much depressed arc departing from the high frequency range, which is well associated with both surface exchange and bulk mass transfer; and hence, one would expect a significant deviation from the symmetrical semicircular shape due to Warburg-like diffusion or what is known as the Gerischer impedance [28,29]. In contrast, for the x = 0.05 sample, the corresponding arc in the complex impedance plane is much close to the ideal semicircle, which suggests that only the

200 400 600

Temperature (°C)

0.000 200

TEC: 16-22 below 900 °C


600 800 Temperature (°C)

Fig. 5. Thermal driven mass loss of Ca3(Coi_xCux)2O6.

Fig. 6. Temperature dependent thermal expansion (AL/Lq) of Ca^Co^ Cux)2O6 measured in air.



fV\ -.-75C


—■— 800 °C

-•- 750 °C

▲ 700 °C

-Y- 650 °C

Table 2

0.0 0.5 1.0 1.5 2.0


(b) 0.2

—■—jc=0.00 —□—jc=0.05

0.1 ■

0.3 Z'(Qcm)

Fig. 7. EIS of Ca3(Coi_xCux)2O6 (x = 0, 0.05): (a) at different temperatures; (b) compared profiles at 800 °C and the proposed equivalent circuit.

surface exchange is retained as the limiting process [29]. Thus the ion transport is accelerated and no longer dominates the rate-determining step. To make these analyses more quantitative, the EIS data were fitted under the proposed equivalent circuit model (inset of Fig. 7b). The obtained values are listed in Table 2. Here, we lay stress on the effect of substitution on the oxygen-ion transport property, and this EIS feature of it is commonly reflected at low frequency, and so the values of Rohm are ignored. We can see that both Rj and R2 of Cu-substituted cathode decrease compared to those of undoped one. So, not only the charge-transfer but also the ionic transport is enhanced as well by Cu substitution. Since parameters R2 and Q2 are more closely associated with ion transport properties, their values are specially used to extract the effective capacitance according to the following expression [30]:

C = (R1-nß)

It turns out that the C2 values reach to 31 and 43 mF cm~ for x = 0 and x = 0.05 samples, respectively. The exceptionally high capacitances further demonstrate the mixed

Fitted results of EIS across Ca3(CoL _xCux)2O6 symmetric electrodes at 800 °C.

Elements x = 0 x = 0.05

Rohm (U) ignored ignored

Q1 (S sec-n cm~2) 0.0248 0.0125

n (Freq. power for Q1) 0.8789 1

R1 (U cm2) 0.0476 0.0135

Q2 (S sec-n cm"2) 0.084 0.112

n (Freq. power for Q2) 0.8066 0.8126

R2 (U cm2) 0.209 0.144

electronic-ionic conducting behavior, as so-called chemical capacitance [28]. It is also pointed out that this chemical capacitance can be viewed as a measure of how much the bulk transport is involved [28], the x = 0.05 sample is therefore endowed with better ionic transport property.

As for the x = 0.1 and x = 0.15 samples, no matter how hard we tried, they could not be fired onto LSGM electrolyte disc under the same sintering condition due to their enlarged TECs, and so their electrochemical performances were not obtained. Actually, since the impurity phase appears in these samples and the surface properties have been changed, the electrochemical data even obtained from them cannot reflect their real performances.

Fig. 8 shows cell voltage and power density as functions of current density for the single cells with the x = 0 and 0.05 samples as cathodes working in a pure H2 flow. A 230-mm-thick LSGM disc was used as electrolyte and Ni-SDC as anode. The Eocv is as high as 1.15 V at 800 °C, very close to the Nernst potential. Maximum power densities (Pmax) for the cell with x = 0 cathode are 406, 340, 230 and 180 mW cm~2, respectively, at 800, 750, 700 and 650 °C (Fig. 8a). These outputs look high as compared with the relatively large ASR values attained from EIS of symmetric electrode, which can be explained by that this type of cathode functions much better under polarization conditions just like Laj_xSrxMnO3 cathode in which bulk path for oxygen-ion transport begins to play an

Fig. 8 single

Current dependent operating potentials as well as power density of cells based on Ca3(Coi_xCux)2O6 cathode at different temperatures.

important role in ORR [28,31]. With Cu substitution, the x = 0.05 cathode exhibits remarkably enhanced power density, as shown in Fig. 8b. At 800 °C, the Pmax reaches as high as 550 mW cm~2. It is worth mentioning that, since our cells operate on thick supported electrolyte, the power density can be further enhanced if the thickness of electrolyte is reduced. Therefore, a low-level Cu substitution in CCO can effectively improve the electrochemical performance, which can be ascribed to the enhanced intrinsic electronic and ionic transport property. The power output in the x = 0.05 cathode agrees well with the above investigated EIS results.

4. Conclusions

Cu substitution is employed to enhance the transport properties of Ca3Co2O6. Studies show that Cu substitution in CCO has great effects on structure, electrical conductivity, thermal expansion and electrochemical performance. Cu substitution leads to the coexistence of both mixed oxidation state of Co and disordered oxygen deficiencies; excessive substitution may bring an impurity phase. As cathode material for SOFC, Ca3(Co1_xCux)2O6 with x = 0.05 exhibits a power density as high as 550 mW cm~2 at 800 °C, which is contributed to the enhanced electronic and ionic transport property. Our experiments indicate that cationic substitution of heteroatom in the Co-based cathode materials is an efficient way to modify the Co valence and oxygen vacancies and hence to improve the electronic/ionic transport and electrochemical performances for SOFCs.


This work was supported by the Natural Science Foundation of China (Grants 513111014 and 21175050). In addition, the authors thank the Analytical and Testing Center of Huazhong University of Science and Technology for XRD, XPS measurements.


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Dr. Fushao Li was born in Guizhou, China in 1978. In June 2001, he graduated from Central South University of China with a B.S. in physicochemisty of metallurgy. In May 2009, he received his Ph.D. degree in Chemical Engineering from Harbin Institute of Technology (China). In October 2012, he came to Huaz-hong University of Science and Technology (China) to undertake post-doctoral research in Professor Yunhui Huang's group. At present, he works in the Qujing Normal University (Yunnan, China). His main research interests include electrode materials of

lithium-ion batteries, cathode catalytic materials for solid-oxide fuel cells, and mechanism investigation of the related electrode processes.

Prof. Yun-Hui Huang received his B.S., M.S. and Ph.D. from Peking University. In 2000, he worked as a postdoctoral researcher in Peking University. From 2002 to 2004, he worked as an associate professor in Fudan University and a JSPS fellow at Tokyo Institute of Technology, Japan. He then worked with Prof. John B. Goodenough in the University of Texas at Austin. In 2008, he became a chair professor of materials science in Huazhong University of Science and Technology. He is now the dean of School of Materials Science and Engineering. His research group work on batteries of energy storage and conversion. For details please see the lab website: