- Email: [email protected]
Sinteractivity, proton conductivity and chemical stability of BaZr0.7In0.3O3-δ for solid oxide fuel cells (SOFCs)
Solid State Ionics 196 (2011) 59–64
Contents lists available at ScienceDirect
Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Sinteractivity, proton conductivity and chemical stability of BaZr0.7In0.3O3-δ for solid oxide fuel cells (SOFCs) Lei Bi, Emiliana Fabbri, Ziqi Sun, Enrico Traversa ⁎ International Center for Materials Nanoarchitectonics (MANA), National Institute for Materials Science (NIMS), 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan
a r t i c l e
i n f o
Article history: Received 20 March 2011 Received in revised form 18 May 2011 Accepted 26 June 2011 Available online 23 July 2011 Keywords: BaZrO3 Ceramic membrane Proton conductor Sintering Solid oxide fuel cell (SOFC)
a b s t r a c t In3+ was used as dopant for BaZrO3 proton conductor and 30 at%-doped BaZrO3 samples (BaZr0.7In0.3O3-δ, BZI) were prepared as electrolyte materials for proton-conducting solid oxide fuel cells (SOFCs). The BZI material showed a much improved sinteractivity compared with the conventional Y-doped BaZrO3. The BZI pellets reached almost full density after sintering at 1600 °C for 10 h, whereas the Y-doped BaZrO3 samples still remained porous under the same sintering conditions. The conductivity measurements indicated that BZI pellets showed smaller bulk but improved grain boundary proton conductivity, when compared with Ydoped BaZrO3 samples. A total proton conductivity of 1.7 × 10 −3 S cm −1 was obtained for the BZI sample at 700 °C in wet 10% H2 atmosphere. The BZI electrolyte material also showed adequate chemical stability against CO2 and H2O, which is promising for application in fuel cells. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Solid oxide fuel cells (SOFCs) are considered as a possible solution to solve the problems of green-house gas emissions because of their low impact to environment, being water the only emission when hydrogen is used as a fuel, and high efficiency as a mean of generating power [1,2]. However, traditional SOFCs require high operating temperatures, which not only lead to high costs but also to longterm stability issues. Therefore, the working temperature reduction is a critical issue for SOFC applications and the development of new materials with high ionic conductivity it is now a hot research topic [3,4]. Iwahara et al. [5] have found that some perovskite oxides show good proton conductivity at intermediate temperatures and thus are promising electrolyte candidates for intermediate-temperature SOFCs. During the past three decades, although many oxides have been found to show certain proton conductivity, the most intensively studied high temperature proton conductors (HTPCs) are doped BaCeO3 and doped BaZrO3 [6]. In polycrystalline pellet form, doped BaCeO3 shows the largest ionic conductivity among the HTPC family [7]. However, its chemical instability in the presence of CO2 and H2O limits its practical use in SOFCs [6,8]. BaCeO3-based materials tend to easily react with CO2 or H2O at intermediate temperatures to form BaCO3 (or Ba(OH)2) and CeO2, which causes electrolyte degradation [8]. Although there are some works concerning the improvement of BaCeO3 chemical stability using different dopants [9–13], these are just improvements and the stability problems for BaCeO3 cannot be
⁎ Corresponding author. Tel.: +81 29 860 4896; fax: +81 29 860 4706. E-mail address: [email protected] (E. Traversa). 0167-2738/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2011.06.014
solved in this way. In comparison, doped BaZrO3 is known to have excellent chemical stability against CO2 and H2O [14–16] as well as high bulk conductivity [17], having a great potential for use as electrolyte material for proton-conducting SOFCs. However, the major problem for Y-doped BaZrO3 (Y is the dopant that allows the best proton conductivity) is its difficult sintering that needs high sintering temperatures (up to 1800 °C) [16]. The high sintering temperature not only leads to high cost but also to the difficulty in choosing a compatible electrode material for electrode-supported fuel cells. Additionally, the poor sinteractivity leads to large grain boundary volume content, which are blocking for proton conductivity, restricting BaZrO3 applications [18]. Thus, the BaZrO3 sinteractivity needs to be improved for its practical applications. The methods proposed to promote the sinteractivity of BaZrO3 are the use of sintering aids (such as ZnO and NiO) [19,20] and the introduction of a second dopant (such as Pr) [21]. Ito et al. [22] have found that the introduction of In 3+ as the second dopant to Y-doped BaZrO3 can lead to slightly better sinterability, indicating that In 3+ can promote the sinteractivity for BaZrO3, and similar results were reported by Imashuku et al.[23] In this study, In 3+ was successfully used as the single dopant for BaZrO3 to improve sinteractivity. Heavily In-doped BaZrO3 (BaZr0.7In0.3O3-δ, BZI) samples showed also an improved grain boundary proton conductivity, which makes BZI a promising electrolyte candidate for proton-conducting SOFCs. 2. Experimental BaZr0.7In0.3O3-δ (BZI) powders were synthesized using a wet chemical route. Stoichiometric amounts of Ba(NO3)2, ZrO(NO3)2·2H2O and In(NO3)3·6H2O were dissolved in distilled water. Citric acid was
60
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
then added as a complexing agent, setting at 1.5 the molar ratio of citric acid with total metal cations. NH4OH was added to the solution to adjust the pH value around 8. The solution was heated under stirring to evaporate water until it changed into a viscous gel and finally ignited to flame, resulting in a white ash. The ash was calcined at different temperatures for 6 h to form the BZI powders. For sake of comparison, BaZr0.8Y0.2O3-δ (BZY) powders were prepared by the same method described above, using Ba(NO3)2, ZrO(NO3)2·2H2O and Y(NO3)3·6H2O as the starting materials. For shrinkage behavior measurements and pellet preparation, BZI and BZY powders heated to 800 °C were used. The shrinkage behavior of a green BZI bar, prepared by uniaxially pressing the powder at 100 MPa for 1 min, was measured from room temperature to 1600 °C using a thermal expansion analyzer (DIL 402E, Netzsch) in air. For comparison, the shrinkage of a green BZY bar was measured using the same procedure. BZI and BZY pellets, 13 mm in diameter and about 1 mm in thickness, were obtained by uniaxially pressing the powders at 200 MPa for 1 min and then sintered at 1600 °C for 10 h in air. X-ray diffraction (XRD, Rigaku, with Cu Ka radiation) analysis was used to identify the phases present in the as-prepared powders and pellets. Scanning electron microscopy (SEM, Hitachi S-4800) was used to observe the morphology of the BZI and BZY powders and pellets after firing at 1600 °C. For conductivity measurements, silver paste electrodes were painted on both sides of the dense BZI pellets and then fired at 700 °C for 2 h. Conductivity measurements were performed in different atmospheres using a multichannel potentiostat (VMP3 Bio-Logic Co.). Anode-supported BZI electrolyte films were fabricated on a NiO-BZI composite anode by a co-pressing and co-sintering procedure. First, the BZI powder was well mixed with NiO in a 40:60 weight ratio to prepare the anode substrates. To enable a sufficient porosity in the anode, 20 wt.% corn flour was added as pore former. Second, a BZI layer was fabricated on the anode by a co-pressing method and the BZI electrolyte thickness was controlled by varying the amounts of the BZI powder, which was heated to 800 °C for 6 h. The anode powder was pressed at 150 MPa into disks 13 mm in diameter and 0.8 mm thick as green substrates. The BZI powder and the substrate were then co-pressed at 200 MPa to form a green bi-layer made of anode substrate and electrolyte. Then, the green bi-layer was co-fired at 1450 °C in air for 6 h to form a half-cell. To test the chemical stability, the bi-layer consisting of a BZI-NiO anode support and a BZI electrolyte layer was treated both in boiling water for 3 h and in pure flowing CO2 environment at 700 °C for 3 h, respectively. The flowing rate of CO2 was set at 250 mL min−1. The samples after treatments were examined by XRD analysis. To assemble a prototype cell, a slurry consisting of PrBaCo2O5 + δ and BaZr0.7Y0.2Pr0.1O3-δ powders mixed in a 1:1 weight ratio was printed on the sintered electrolyte membrane surface and then fired at 1000 °C for 3 h to form a porous cathode. A single cell with a NiO-BZI (anode)/ BZI (electrolyte)/ PrBaCo2O5 + δ-BaZr0.7Y0.2Pr0.1O3-δ (cathode) structure was assembled. Humidified hydrogen (~3% H2O) was fed to the anode chamber at a flow rate of 25 mL min−1, while the cathode was exposed to atmospheric air. The anode side was sealed with Ag paste. Two silver wires were connected to each electrode as current collectors. The performance of the fuel cell was measured at different temperatures using a multichannel potentiostat (VMP3 Bio-Logic Co.). Electrochemical impedance spectroscopy (EIS) measurements were performed under open circuit voltage ranging the frequency from 0.1 Hz to 1 MHz, with an AC voltage amplitude of 100 mV. The morphology of the tested cells was observed using a Hitachi S-4800 SEM.
Fig. 1. XRD patterns for (a) BZI powders fired at 800 °C and 1100 °C, (b) BZI pellet sintered at 1600 °C for 10 h. The reflection lines of undoped BaZrO3 (JCPDS Card No. 060399) are also shown as reference.
firing at 800 °C for 6 h, even though some BaCO3 peaks were also present. When the firing temperature was increased up to 1100 °C, only the BZI phase peaks were found. The formation of the BaZr0.7In0.3O3-δ single phase after firing at 1100 °C is in agreement with previous reports that showed In 3+ doping concentration reaching 50 at.% for BaZrO3 [24,25]. The BZI powder calcined at 800 °C, despite the residual presence of BaCO3, was pressed into pellets and sintered at 1600 °C for 10 h for conductivity tests, since Yamazaki et al. [26] reported that the incomplete calcination was helpful to improve both sinteractivity and proton conductivity of Y-doped BaZrO3. Fig. 2(b) shows the XRD pattern of the BZI pellet after sintering at 1600 °C, which indicates that no secondary phases could be observed.
3. Results and discussion 3.1. XRD analysis Fig. 1(a) shows the XRD patterns of the BZI powders fired at different temperatures. The main perovskite phase was formed after
Fig. 2. Shrinkage of green BZI and BZY bars tested from room temperature to 1600 °C.
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
61
Fig. 3. SEM micrographs of surface (a) and fracture (b) of BZI pellet after sintering at 1600 °C, and of surface (c) and fracture (d) of BZY pellet after sintering at 1600 °C.
3.2. Sinteractivity Fig. 2 shows the shrinkage behavior of the BZI and conventional BZY samples. Both the green electrolytes were heated from room temperature to 1600 °C with a heating rate of 10 °C min −1. Both samples began to shrink at about 1200 °C. However, the BZI sintering was accelerated with respect to that of BZY. The total shrinkage of the BZI sample was 26.6%, which almost doubled the shrinkage of the BZY sample (~ 13.7%). Obviously, the heavily In-doped BaZrO3 sample shows much better sinteractivity than the conventional Y-doped sample. The improvement in sinteractivity for In-doped BaZrO3 may be due to the low melting point of In2O3 [22]. Fig. 3 shows the SEM micrographs of the BZI and BZY pellets after sintering at 1600 °C for 10 h. The surface (Fig. 3a) and fracture (Fig. 3b) micrographs of the BZI pellet after sintering showed that the
Fig. 4. Complex-impedance plane plots for BZI pellet at 400 °C and 600 °C in wet air atmosphere.
BZI pellet was quite dense without any pores. The surface (Fig. 3c) and fracture (Fig. 3d) SEM micrographs of the BZY pellet prepared with the same procedure after sintering showed that the BZY pellet was still porous. The comparison between the dense BZI and porous BZY samples, obtained after sintering at the same conditions, indicates that 30 at.% In doping can dramatically improve the BaZrO3 sinteractivity, showing promise for its use as SOFC electrolyte. 3.3. Conductivity In addition to good sinteractivity, an electrolyte material should have a suitable ionic conductivity to reduce the overall cell resistance and improve the performance. Fig. 4 shows the complex impedance plane plots of the BZI pellet measured at 400 °C and 600 °C. At 400 °C, the EIS plot showed two depressed arcs. The first arc (labeled as arc 1) appears in the high frequency region with a capacitance in the order of 10 −9 F cm −1, which is typical of grain boundary conductivity [27]. The second arc (labeled as arc 2) in the intermediate frequency range is associated with a capacitance in the order of 10 −6 F cm −1, which is
Fig. 5. Temperature-dependence of the BZI conductivity in wet 10% H2, in dry air and wet air; p(H2O) = 3 kPa.
62
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
Table 1 Comparison of bulk, grain boundary, and total conductivities of the BZI in this study and reported Y-doped BaZrO3 in literatures at 400 °C. Material [reference]
Bulk (S cm−1)
Grain boundary (S cm−1)
Total (S cm−1)
BZY20 [30] BZY-Zn [32] BZI [present study]
1.3 × 10−3 3.3 × 10−4 3.1 × 10−4
6.4 × 10−4 6.9 × 10−4 1.5 × 10−3
4.3 × 10−4 2.2 × 10−4 2.6 × 10−4
related to processes at the electrode/electrolyte interface [28]. The contribution of bulk and grain boundary conductivities to total conductivity can be separated at this temperature. However, when the testing temperature increased to 600 °C, the arc 1 disappeared and only the arc 2 with a capacitance of about 10 −6 F cm −1 could be seen in the EIS plot. Fig. 5 shows the Arrhenius plots of the conductivity values for the BZI pellet, obtained from EIS measurements in different atmospheres. Below 500 °C, the conductivities of the sample in wet 10% H2 and in wet air were close and much larger than the conductivity in dry air. This indicated that the main charge carriers in wet atmosphere were protons, formed according to the following reaction: ••
×
•
H2 O + Vo + Oo ⇔2OH :
ð1Þ
The activation energy values in wet 10% H2 and wet air were 0.45 eV and 0.66 eV, respectively, in agreement with the literature [15,16]. At higher temperatures, the conductivity in wet air was larger than that in wet 10% H2 due to the additional contribution of electron hole conduction in an oxidative atmosphere, described in the following reaction: ••
Vo +
1 × • O ⇔Oo + h : 2 2
ð2Þ
The activation energy of BZI conductivity in dry air was 0.87 eV, indicating that in dry atmosphere the charge carriers were oxygen ions and electron holes, instead of protons, according to (2). Above 600 °C, the total conductivity of the sample in dry air was larger than those in wet atmospheres, which can be explained by the competition between (1) and (2); with the presence of water, the formation of proton carriers, according to (1), partially diminished the number of oxygen ions and electron holes. The total conductivity of BZI in wet 10% H2 atmosphere (balanced with Ar) reached 1.7 × 10 −3 S cm −1 at 700 °C, which is smaller than the total conductivity values reported for well sintered Y-doped BaZrO3 materials in wet H2 atmosphere [15,16,26], although the reported BZY conductivity varies significantly due to different preparation processes [6,29]. Table 1 reports the BZI bulk and grain boundary conductivities, obtained from the EIS plot at 400 °C, to have
Fig. 6. XRD patterns of an anode supported BZI membrane fired at 1450 °C before (a), and after exposure to pure CO2 at 700 °C for 3 h (b), and to boiling water for 3 h (c).
Fig. 7. XRD patterns of a composite NiO-BZI anode fired at 1450 °C before (a), and after exposure to pure CO2 at 700 °C for 3 h (b), and to boiling water for 3 h (c).
a better understanding of the In-doping influence on the BaZrO3 conductivity. Compared with 20% Y-doped BaZrO3 (BZY20) [30], BZI shows a smaller bulk conductivity, as expected. However, it is interesting to observe that the BZI grain boundary conductivity is much improved with respect to BZY. The grain boundary conductivity is related to the grain boundary volume content and fewer boundaries will be encountered in the process of proton transport for large grainsized samples [31]. Therefore, the larger grain boundary conductivity for BZI can be attributed to the improved sinteractivity that increased grain size up to about 3 μm, effectively reducing the grain boundary volume content. Even compared with the ZnO-modified BZY (BZY-Zn) [32], the BZI sample shows better grain boundary and total conductivities (Table 1). The large BZI grain boundary conductivity, added to its good sinteractivity, makes this material interesting for BaZrO3-based fuel cell applications. The smaller conductivity than the conventional Ydoped BaZrO3 may be compensated by the better processability that may allow fabricating electrode-supported electrolyte film configuration: using tens of microns thick electrolyte films can reduce the ohmic resistance and compensate the negative effect of reduced conductivity [33]. 3.4. Chemical stability Good chemical stability is a crucial feature for SOFC electrolyte materials, especially operating with carbon-containing fuels. This issue is even more important for electrolyte films due to the reduced thickness [6,8]. Therefore, the chemical stability of BZI electrolyte layer on NiO-BZI anode half-cell was investigated. Fig. 6 shows the XRD patterns of the as-prepared BZI films and after exposure to pure
Fig. 8. Electrochemical performance (I-V and power density curves) of a BZI-based fuel cell measured from 500 to 650 °C.
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
63
Fig. 10. Typical complex-impedance plane plot of the BZI single cell measured at 600 °C at open circuit voltage.
Fig. 9. SEM cross-sectional micrograph of the BZI single cell after testing.
CO2 and boiling water. The XRD pattern of the sintered BZI (Fig. 6a) confirms that the BZI film after sintering at 1450 °C for 6 h was made of a single phase. The XRD pattern of BZI remained unchanged after CO2 exposure at 700 °C for 3 h (Fig. 6b) and after treatment in boiling water for 3 h (Fig. 6c), demonstrating a sufficient chemical stability against CO2 and against water. The chemical stability of the BZI-NiO anode was also investigated since the anode working condition can be aggressive when carboncontaining fuels are used. Moreover, stronger reactivity with CO2 and H2O may occur due to the anode porosity. Fig. 7 shows the XRD patterns of the NiO-BZI composite anode before (Fig. 7a) and after exposure to pure CO2 at 700 °C for 3 h (Fig. 7b), as well as after treatment in boiling water (Fig. 7c). Also in this case the BZI-NiO showed a good chemical stability, since no other peaks except for those of BZI and NiO could be observed. 3.5. Fuel cell performance The greatest advantage of BaZrO3-based materials for SOFC applications is their excellent chemical stability. However, the performance of BaZrO3-based fuel cells is poor, up to now. The difficult processing of Y-doped BaZrO3 allowed mostly the fabrication of electrolyte-supported fuel cells, which achieved several mW cm −2 at elevated temperatures (such as 700 °C) [34] as thick electrolyte layer (such as 1 mm) led to a large ohmic resistance for the cells, which inevitably resulted in poor cell performance. The good sinteractivity of BZI allowed fabricating anode-supported cells with reduced electrolyte thickness to tens of μm, leading to an ohmic cell resistance much smaller than that of electrolyte supported BZY cell [15,21]. Fig. 8 shows the I-V and power density curves at different temperatures for a BZI-based half-cell with a composite cathode made of PrBaCo2O5 + δ and BaZr0.7Y0.2Pr0.1O3-δ. The open circuit voltage (OCV)
was recorded at 0.985, 0.966, 0.946 and 0.929 V at 500, 550, 600 and 650 °C, respectively. These values are close to the OCV values reported in the literature for BaZrO3 electrolyte film cells [21], suggesting that the BZI electrolyte membrane was dense. With humidified hydrogen (~3% H2O) as a fuel and static air as an oxidant, the maximum power densities were 95, 84, 70 and 56 mW cm−2 at 650, 600, 550 and 500 °C, respectively. Fig. 9 shows the SEM micrograph of the single cell cross-section after testing. The BZI electrolyte layer was about 15 μm in thickness and rather dense, without any obvious pores or cracks, agreeing with the measured OCV values, as discussed above. The BZI electrolyte adhered well with both the anode and cathode layers. Table 2 compares the fuel cell performance the BZI-based cell with that of BaZrO3-based fuel cells reported in the relevant literature. Most studies for BaZrO3-based cells were performed using the electrolyte-supported configuration, and the performance of these cells was poor. Only a few studies concerning cells having BaZrO3 electrolyte films were reported and the cell performance was improved compared with that of the electrolyte-supported cells. The present BZI-based cell showed a power output of 84 mW cm −2 at 600 °C, much larger than that obtained for electrolyte-supported cells, as expected, but comparable to that of electrode-supported BZY cells. The improved sinteractivity, which enhanced proton transport at the grain boundaries and allowed sintering at lower temperatures, is the reason for the good fuel cell performance together with the acceptable conductivity of BZI. Although the current BZI cell showed larger power density than those of the most fuel cells listed in the table, a slightly larger power density output of 110 mW cm −2 at 600 °C was achieved using a Y-doped BaZrO3 4 μm-thick electrolyte film prepared by pulsed laser deposition (PLD) [40]. PLD is more suitable for smallscale applications [41], and a further reduction in the BZI electrolyte thickness may be beneficial to enhance the cell performance. Fuel cell performance is not only affected by the conductivity and thickness of the electrolyte but also by the electrode polarization, which increases with reducing the operation temperature [42]. EIS measurements at open circuit conditions of the fuel cell were used to evaluate the different resistances in the cell. Fig. 10 shows a typical EIS plot of
Table 2 Comparison of the performance of stable BaZrO3-based fuel cells reported in the literature, and the BZI-based cell developed in the present study. LSM: (La0.8Sr0.2)0.98MnO3-δ; BSCF: Ba0.5Sr0.5Co0.8Fe0.2O3-δ; LSCF:La0.6Sr0.4Co0.2Fe0.8O3-δ; PBCO:PrBaCo2O5 + δ. Year [reference]
Cell configuration
Electrolyte thickness (μm)
Cathode / Anode
Peak power density (mW cm−2)
2008 [34] 2008 [35] 2008 [15] 2009 [36] 2010 [37] 2010 [38] 2009 [36] 2009 [39] 2010 [40] Present work
Electrolyte supported Electrolyte supported Electrolyte supported Electrolyte supported Electrolyte supported Electrolyte supported Cathode supported Anode supported Anode supported Anode supported
1000 1000 600 1000 1200 300 10 30 4 15
Pt / Pt Pt / Pt Pt / Pt Pt / Pt Pt / Pt Pt / Pt LSM / NiO BSCF / NiO LSCF / NiO PBCO / NiO
7 (at 700 °C) 4 (at 700 °C) 7 (at 700 °C) 4 (at 800 °C) 3 (at 800 °C) 11 (at 800 °C) 26 (at 800 °C) 23 (at 600 °C) 110 (at 600 °C) 84 (at 600 °C)
64
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
Acknowledgments This work was supported in part by the World Premier International Research Center Initiative of MEXT, Japan. The authors thank Dr. Hidehiko Tanaka for his invaluable technical assistance.
References [1] [2] [3] [4] [5] [6] [7] Fig. 11. Temperature-dependence of the BZI-based cell total, ohmic and interfacial polarization resistances, and of the ratio between ohmic resistance and cell total resistance.
the cell, measured at 600 °C. The intercept with the real axis at high frequency represents the ohmic resistance of the cell (Rohmic), which includes the electrolyte resistance and lead wires. The low frequency intercept corresponds to the total resistance of the cell. Therefore, the difference between the high frequency and low frequency intercepts with the real axis represents the total interfacial polarization resistance (Rp) of the cell. Fig. 11 shows the temperature-dependence of the total resistance, Rohmic and Rp of the cell, as determined from the EIS data. The Rohmic of the cell was 2.82, 2.39, 2.01 and 1.83 Ω cm2 at 500, 550, 600 and 650 °C, respectively, whereas the Rp values were 1.41, 0.57, 0.27 and 0.16 Ω cm2 at 500, 550, 600 and 650 °C, respectively. The ratio between Rohmic and the total resistance increased from 66% at 500 °C to 92% at 650 °C, indicating that the Rohmic makes a larger contribution to the cell total resistance in the whole testing temperature range, becoming dominant at high temperatures. Although the lower total conductivity of BZI compared with conventional Y-doped BaZrO3 partially increased Rohmic, larger electrolyte thickness and the interfacial resistances could also be the reason for the large Rohmic [11,40]. Therefore, further work, including surface modification, fabrication of electrode functional layer, and further reduction of electrolyte thickness, needs to be carried out to further improve the BZI cell performance. 4. Conclusions BaZr0.7In0.3O3-δ (BZI) was synthesized for protonic-SOFCs to overcome the poor sinteractivity of Y-doped BaZrO3, without detrimental effects on the good chemical stability. The BZI sinteractivity was improved and also the grain boundary conductivity was increased, but the bulk conductivity was lowered. Despite the total electrolyte conductivity was reduced, anode-supported fuel cell with a 15 μm thick BZI electrolyte showed a promising performance, as compared to BaZrO3-based fuel cells. The good chemical stability, improved sinteractivity and promising fuel cell performance suggested that BZI might be a good electrolyte candidate for proton-conducting SOFCs.
[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42]
S.M. Haile, Acta Mater. 51 (2003) 5981. A. Boudghene Stambouli, E. Traversa, Renew. Sustain. Energy Rev. 6 (2002) 433. A. Bieberle-Hutter, J.L. Hertz, H.L. Tuller, Acta Mater. 56 (2008) 177. D.J.L. Brett, A. Atkinson, N.P. Brandon, S.J. Skinner, Chem. Soc. Rev. 37 (2008) 1568. H. Iwahara, H. Uchida, K. Ono, K. Ogaki, J. Electrochem. Soc. 135 (1988) 529. E. Fabbri, D. Pergolesi, E. Traversa, Chem. Soc. Rev. 39 (2010) 4355. T. Hibino, A. Hashimoto, M. Suzuki, M. Sano, J. Electrochem. Soc. 149 (2002) A1503. L. Bi, Z.T. Tao, R.R. Peng, W. Liu, J. Inorg. Mater. 25 (2010) 1. Z.M. Zhong, Solid State Ionics 178 (2007) 213. C.D. Zuo, S.W. Zha, M.L. Liu, M. Hatano, M. Uchiyama, Adv. Mater. 18 (2006) 3318. L. Bi, S.Q. Zhang, S.M. Fang, Z.T. Tao, R.R. Peng, W. Liu, Electrochem. Commun. 10 (2008) 1598. L. Bi, Z.T. Tao, C. Liu, W.P. Sun, H.Q. Wang, W. Liu, J. Membr. Sci. 336 (2009) 1. A. Tomita, K. Tsunekawa, T. Hibino, S. Teranishi, Y. Tachi, M. Sano, Solid State Ionics 177 (2006) 2951. H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, H. Suzuki, Solid State Ionics 61 (1993) 65. E. Fabbri, A. D'Epifanio, E. Di Bartolomeo, S. Licoccia, E. Traversa, Solid State Ionics 179 (2008) 558. K. Katahira, Y. Kohchi, T. Shimura, H. Iwahara, Solid State Ionics 138 (2000) 91. D. Pergolesi, E. Fabbri, A. D'Epifanio, E. Di Bartolomeo, A. Tebano, S. Sanna, S. Licoccia, G. Balestrino, E. Traversa, Nat. Mater. 9 (2010) 846. P. Babilo, S.M. Haile, J. Am. Ceram. Soc. 88 (2005) 2362. S.W. Tao, J.T.S. Irvine, Adv. Mater. 18 (2006) 1581. J.H. Tong, D. Clark, M. Hoban, R. O' Hayre, Solid State Ionics 181 (2010) 496. E. Fabbri, L. Bi, H. Tanaka, D. Pergolesi, E. Traversa, Adv. Funct. Mater. 21 (2011) 158. N. Ito, H. Matsumoto, Y. Kawasaki, S. Okada, T. Ishihara, Solid State Ionics 179 (2008) 324. S. Imashuku, T. Uda, Y. Nose, G. Taniguchi, Y. Ito, Y. Awakura, J. Electrochem. Soc. 156 (2009) B1. F. Giannici, A. Longo, A. Balerna, K.D. Kreuer, A. Martorana, Chem. Mater. 21 (2009) 2641. I. Ahmed, S.G. Eriksson, E. Ahlberg, C.S. Knee, Solid State Ionics 179 (2008) 1155. Y. Yamazaki, R. Hernandez-Sanchez, S.M. Haile, Chem. Mater. 21 (2009) 2755. T.S. Bjorheim, A. Kuwabara, I. Ahmed, R. Haugsrud, S. Stolen, T. Norby, Solid State Ionics 181 (2010) 130. V. Dusastre, J.A. Kilner, Solid State Ionics 126 (1999) 163. S.B.C. Duval, P. Holtappels, U.F. Vogt, U. Stimming, T. Graule, Fuel Cells 5 (2009) 613. E. Fabbri, D. Pergolesi, S. Licoccia, E. Traversa, Solid State Ionics 181 (2010) 1043. S.M. Haile, G. Staneff, K.H. Ryu, J. Mater. Sci. 36 (2001) 1149. S.W. Tao, J.T.S. Irvine, J. Solid State Chem. 180 (2007) 3493. N.H. Menzler, F. Tietz, S. Uhlenbruck, H.P. Buchkremer, D. Stover, J. Mater. Sci. 45 (2010) 3109. A. D'Epifanio, E. Fabbri, E. Di Bartolomeo, S. Licoccia, E. Traversa, Fuel Cells 8 (2008) 69. E. Fabbri, D. Pergolesi, A. D'Epifanio, E. Di Bartolomeo, G. Balestrino, S. Licoccia, E. Traversa, Energy Environ. Sci. 1 (2008) 355. C. Peng, J. Melnik, J.X. Li, J.L. Luo, A.R. Sanger, K.T. Chuang, J. Power Sour. 190 (2009) 447. W.M. Guo, J.A. Liu, C. Jin, J. Alloys Compd. 504 (2010) L21. A.L. Tsai, M. Kopczyk, R.J. Smith, V.H. Schmidt, Solid State Ionics 181 (2010) 1083. Y.M. Guo, Y. Lin, R. Ran, Z.P. Shao, J. Power Sour. 193 (2009) 400. D. Pergolesi, E. Fabbri, E. Traversa, Electrochem. Commun. 12 (2010) 977. E. Traversa, Interface 18 (2009) 49. E. Fabbri, D. Pergolesi, E. Traversa, Sci. Technol. Adv. Mater. 11 (2010) 044301.
Contents lists available at ScienceDirect
Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Sinteractivity, proton conductivity and chemical stability of BaZr0.7In0.3O3-δ for solid oxide fuel cells (SOFCs) Lei Bi, Emiliana Fabbri, Ziqi Sun, Enrico Traversa ⁎ International Center for Materials Nanoarchitectonics (MANA), National Institute for Materials Science (NIMS), 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan
a r t i c l e
i n f o
Article history: Received 20 March 2011 Received in revised form 18 May 2011 Accepted 26 June 2011 Available online 23 July 2011 Keywords: BaZrO3 Ceramic membrane Proton conductor Sintering Solid oxide fuel cell (SOFC)
a b s t r a c t In3+ was used as dopant for BaZrO3 proton conductor and 30 at%-doped BaZrO3 samples (BaZr0.7In0.3O3-δ, BZI) were prepared as electrolyte materials for proton-conducting solid oxide fuel cells (SOFCs). The BZI material showed a much improved sinteractivity compared with the conventional Y-doped BaZrO3. The BZI pellets reached almost full density after sintering at 1600 °C for 10 h, whereas the Y-doped BaZrO3 samples still remained porous under the same sintering conditions. The conductivity measurements indicated that BZI pellets showed smaller bulk but improved grain boundary proton conductivity, when compared with Ydoped BaZrO3 samples. A total proton conductivity of 1.7 × 10 −3 S cm −1 was obtained for the BZI sample at 700 °C in wet 10% H2 atmosphere. The BZI electrolyte material also showed adequate chemical stability against CO2 and H2O, which is promising for application in fuel cells. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Solid oxide fuel cells (SOFCs) are considered as a possible solution to solve the problems of green-house gas emissions because of their low impact to environment, being water the only emission when hydrogen is used as a fuel, and high efficiency as a mean of generating power [1,2]. However, traditional SOFCs require high operating temperatures, which not only lead to high costs but also to longterm stability issues. Therefore, the working temperature reduction is a critical issue for SOFC applications and the development of new materials with high ionic conductivity it is now a hot research topic [3,4]. Iwahara et al. [5] have found that some perovskite oxides show good proton conductivity at intermediate temperatures and thus are promising electrolyte candidates for intermediate-temperature SOFCs. During the past three decades, although many oxides have been found to show certain proton conductivity, the most intensively studied high temperature proton conductors (HTPCs) are doped BaCeO3 and doped BaZrO3 [6]. In polycrystalline pellet form, doped BaCeO3 shows the largest ionic conductivity among the HTPC family [7]. However, its chemical instability in the presence of CO2 and H2O limits its practical use in SOFCs [6,8]. BaCeO3-based materials tend to easily react with CO2 or H2O at intermediate temperatures to form BaCO3 (or Ba(OH)2) and CeO2, which causes electrolyte degradation [8]. Although there are some works concerning the improvement of BaCeO3 chemical stability using different dopants [9–13], these are just improvements and the stability problems for BaCeO3 cannot be
⁎ Corresponding author. Tel.: +81 29 860 4896; fax: +81 29 860 4706. E-mail address: [email protected] (E. Traversa). 0167-2738/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2011.06.014
solved in this way. In comparison, doped BaZrO3 is known to have excellent chemical stability against CO2 and H2O [14–16] as well as high bulk conductivity [17], having a great potential for use as electrolyte material for proton-conducting SOFCs. However, the major problem for Y-doped BaZrO3 (Y is the dopant that allows the best proton conductivity) is its difficult sintering that needs high sintering temperatures (up to 1800 °C) [16]. The high sintering temperature not only leads to high cost but also to the difficulty in choosing a compatible electrode material for electrode-supported fuel cells. Additionally, the poor sinteractivity leads to large grain boundary volume content, which are blocking for proton conductivity, restricting BaZrO3 applications [18]. Thus, the BaZrO3 sinteractivity needs to be improved for its practical applications. The methods proposed to promote the sinteractivity of BaZrO3 are the use of sintering aids (such as ZnO and NiO) [19,20] and the introduction of a second dopant (such as Pr) [21]. Ito et al. [22] have found that the introduction of In 3+ as the second dopant to Y-doped BaZrO3 can lead to slightly better sinterability, indicating that In 3+ can promote the sinteractivity for BaZrO3, and similar results were reported by Imashuku et al.[23] In this study, In 3+ was successfully used as the single dopant for BaZrO3 to improve sinteractivity. Heavily In-doped BaZrO3 (BaZr0.7In0.3O3-δ, BZI) samples showed also an improved grain boundary proton conductivity, which makes BZI a promising electrolyte candidate for proton-conducting SOFCs. 2. Experimental BaZr0.7In0.3O3-δ (BZI) powders were synthesized using a wet chemical route. Stoichiometric amounts of Ba(NO3)2, ZrO(NO3)2·2H2O and In(NO3)3·6H2O were dissolved in distilled water. Citric acid was
60
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
then added as a complexing agent, setting at 1.5 the molar ratio of citric acid with total metal cations. NH4OH was added to the solution to adjust the pH value around 8. The solution was heated under stirring to evaporate water until it changed into a viscous gel and finally ignited to flame, resulting in a white ash. The ash was calcined at different temperatures for 6 h to form the BZI powders. For sake of comparison, BaZr0.8Y0.2O3-δ (BZY) powders were prepared by the same method described above, using Ba(NO3)2, ZrO(NO3)2·2H2O and Y(NO3)3·6H2O as the starting materials. For shrinkage behavior measurements and pellet preparation, BZI and BZY powders heated to 800 °C were used. The shrinkage behavior of a green BZI bar, prepared by uniaxially pressing the powder at 100 MPa for 1 min, was measured from room temperature to 1600 °C using a thermal expansion analyzer (DIL 402E, Netzsch) in air. For comparison, the shrinkage of a green BZY bar was measured using the same procedure. BZI and BZY pellets, 13 mm in diameter and about 1 mm in thickness, were obtained by uniaxially pressing the powders at 200 MPa for 1 min and then sintered at 1600 °C for 10 h in air. X-ray diffraction (XRD, Rigaku, with Cu Ka radiation) analysis was used to identify the phases present in the as-prepared powders and pellets. Scanning electron microscopy (SEM, Hitachi S-4800) was used to observe the morphology of the BZI and BZY powders and pellets after firing at 1600 °C. For conductivity measurements, silver paste electrodes were painted on both sides of the dense BZI pellets and then fired at 700 °C for 2 h. Conductivity measurements were performed in different atmospheres using a multichannel potentiostat (VMP3 Bio-Logic Co.). Anode-supported BZI electrolyte films were fabricated on a NiO-BZI composite anode by a co-pressing and co-sintering procedure. First, the BZI powder was well mixed with NiO in a 40:60 weight ratio to prepare the anode substrates. To enable a sufficient porosity in the anode, 20 wt.% corn flour was added as pore former. Second, a BZI layer was fabricated on the anode by a co-pressing method and the BZI electrolyte thickness was controlled by varying the amounts of the BZI powder, which was heated to 800 °C for 6 h. The anode powder was pressed at 150 MPa into disks 13 mm in diameter and 0.8 mm thick as green substrates. The BZI powder and the substrate were then co-pressed at 200 MPa to form a green bi-layer made of anode substrate and electrolyte. Then, the green bi-layer was co-fired at 1450 °C in air for 6 h to form a half-cell. To test the chemical stability, the bi-layer consisting of a BZI-NiO anode support and a BZI electrolyte layer was treated both in boiling water for 3 h and in pure flowing CO2 environment at 700 °C for 3 h, respectively. The flowing rate of CO2 was set at 250 mL min−1. The samples after treatments were examined by XRD analysis. To assemble a prototype cell, a slurry consisting of PrBaCo2O5 + δ and BaZr0.7Y0.2Pr0.1O3-δ powders mixed in a 1:1 weight ratio was printed on the sintered electrolyte membrane surface and then fired at 1000 °C for 3 h to form a porous cathode. A single cell with a NiO-BZI (anode)/ BZI (electrolyte)/ PrBaCo2O5 + δ-BaZr0.7Y0.2Pr0.1O3-δ (cathode) structure was assembled. Humidified hydrogen (~3% H2O) was fed to the anode chamber at a flow rate of 25 mL min−1, while the cathode was exposed to atmospheric air. The anode side was sealed with Ag paste. Two silver wires were connected to each electrode as current collectors. The performance of the fuel cell was measured at different temperatures using a multichannel potentiostat (VMP3 Bio-Logic Co.). Electrochemical impedance spectroscopy (EIS) measurements were performed under open circuit voltage ranging the frequency from 0.1 Hz to 1 MHz, with an AC voltage amplitude of 100 mV. The morphology of the tested cells was observed using a Hitachi S-4800 SEM.
Fig. 1. XRD patterns for (a) BZI powders fired at 800 °C and 1100 °C, (b) BZI pellet sintered at 1600 °C for 10 h. The reflection lines of undoped BaZrO3 (JCPDS Card No. 060399) are also shown as reference.
firing at 800 °C for 6 h, even though some BaCO3 peaks were also present. When the firing temperature was increased up to 1100 °C, only the BZI phase peaks were found. The formation of the BaZr0.7In0.3O3-δ single phase after firing at 1100 °C is in agreement with previous reports that showed In 3+ doping concentration reaching 50 at.% for BaZrO3 [24,25]. The BZI powder calcined at 800 °C, despite the residual presence of BaCO3, was pressed into pellets and sintered at 1600 °C for 10 h for conductivity tests, since Yamazaki et al. [26] reported that the incomplete calcination was helpful to improve both sinteractivity and proton conductivity of Y-doped BaZrO3. Fig. 2(b) shows the XRD pattern of the BZI pellet after sintering at 1600 °C, which indicates that no secondary phases could be observed.
3. Results and discussion 3.1. XRD analysis Fig. 1(a) shows the XRD patterns of the BZI powders fired at different temperatures. The main perovskite phase was formed after
Fig. 2. Shrinkage of green BZI and BZY bars tested from room temperature to 1600 °C.
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
61
Fig. 3. SEM micrographs of surface (a) and fracture (b) of BZI pellet after sintering at 1600 °C, and of surface (c) and fracture (d) of BZY pellet after sintering at 1600 °C.
3.2. Sinteractivity Fig. 2 shows the shrinkage behavior of the BZI and conventional BZY samples. Both the green electrolytes were heated from room temperature to 1600 °C with a heating rate of 10 °C min −1. Both samples began to shrink at about 1200 °C. However, the BZI sintering was accelerated with respect to that of BZY. The total shrinkage of the BZI sample was 26.6%, which almost doubled the shrinkage of the BZY sample (~ 13.7%). Obviously, the heavily In-doped BaZrO3 sample shows much better sinteractivity than the conventional Y-doped sample. The improvement in sinteractivity for In-doped BaZrO3 may be due to the low melting point of In2O3 [22]. Fig. 3 shows the SEM micrographs of the BZI and BZY pellets after sintering at 1600 °C for 10 h. The surface (Fig. 3a) and fracture (Fig. 3b) micrographs of the BZI pellet after sintering showed that the
Fig. 4. Complex-impedance plane plots for BZI pellet at 400 °C and 600 °C in wet air atmosphere.
BZI pellet was quite dense without any pores. The surface (Fig. 3c) and fracture (Fig. 3d) SEM micrographs of the BZY pellet prepared with the same procedure after sintering showed that the BZY pellet was still porous. The comparison between the dense BZI and porous BZY samples, obtained after sintering at the same conditions, indicates that 30 at.% In doping can dramatically improve the BaZrO3 sinteractivity, showing promise for its use as SOFC electrolyte. 3.3. Conductivity In addition to good sinteractivity, an electrolyte material should have a suitable ionic conductivity to reduce the overall cell resistance and improve the performance. Fig. 4 shows the complex impedance plane plots of the BZI pellet measured at 400 °C and 600 °C. At 400 °C, the EIS plot showed two depressed arcs. The first arc (labeled as arc 1) appears in the high frequency region with a capacitance in the order of 10 −9 F cm −1, which is typical of grain boundary conductivity [27]. The second arc (labeled as arc 2) in the intermediate frequency range is associated with a capacitance in the order of 10 −6 F cm −1, which is
Fig. 5. Temperature-dependence of the BZI conductivity in wet 10% H2, in dry air and wet air; p(H2O) = 3 kPa.
62
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
Table 1 Comparison of bulk, grain boundary, and total conductivities of the BZI in this study and reported Y-doped BaZrO3 in literatures at 400 °C. Material [reference]
Bulk (S cm−1)
Grain boundary (S cm−1)
Total (S cm−1)
BZY20 [30] BZY-Zn [32] BZI [present study]
1.3 × 10−3 3.3 × 10−4 3.1 × 10−4
6.4 × 10−4 6.9 × 10−4 1.5 × 10−3
4.3 × 10−4 2.2 × 10−4 2.6 × 10−4
related to processes at the electrode/electrolyte interface [28]. The contribution of bulk and grain boundary conductivities to total conductivity can be separated at this temperature. However, when the testing temperature increased to 600 °C, the arc 1 disappeared and only the arc 2 with a capacitance of about 10 −6 F cm −1 could be seen in the EIS plot. Fig. 5 shows the Arrhenius plots of the conductivity values for the BZI pellet, obtained from EIS measurements in different atmospheres. Below 500 °C, the conductivities of the sample in wet 10% H2 and in wet air were close and much larger than the conductivity in dry air. This indicated that the main charge carriers in wet atmosphere were protons, formed according to the following reaction: ••
×
•
H2 O + Vo + Oo ⇔2OH :
ð1Þ
The activation energy values in wet 10% H2 and wet air were 0.45 eV and 0.66 eV, respectively, in agreement with the literature [15,16]. At higher temperatures, the conductivity in wet air was larger than that in wet 10% H2 due to the additional contribution of electron hole conduction in an oxidative atmosphere, described in the following reaction: ••
Vo +
1 × • O ⇔Oo + h : 2 2
ð2Þ
The activation energy of BZI conductivity in dry air was 0.87 eV, indicating that in dry atmosphere the charge carriers were oxygen ions and electron holes, instead of protons, according to (2). Above 600 °C, the total conductivity of the sample in dry air was larger than those in wet atmospheres, which can be explained by the competition between (1) and (2); with the presence of water, the formation of proton carriers, according to (1), partially diminished the number of oxygen ions and electron holes. The total conductivity of BZI in wet 10% H2 atmosphere (balanced with Ar) reached 1.7 × 10 −3 S cm −1 at 700 °C, which is smaller than the total conductivity values reported for well sintered Y-doped BaZrO3 materials in wet H2 atmosphere [15,16,26], although the reported BZY conductivity varies significantly due to different preparation processes [6,29]. Table 1 reports the BZI bulk and grain boundary conductivities, obtained from the EIS plot at 400 °C, to have
Fig. 6. XRD patterns of an anode supported BZI membrane fired at 1450 °C before (a), and after exposure to pure CO2 at 700 °C for 3 h (b), and to boiling water for 3 h (c).
Fig. 7. XRD patterns of a composite NiO-BZI anode fired at 1450 °C before (a), and after exposure to pure CO2 at 700 °C for 3 h (b), and to boiling water for 3 h (c).
a better understanding of the In-doping influence on the BaZrO3 conductivity. Compared with 20% Y-doped BaZrO3 (BZY20) [30], BZI shows a smaller bulk conductivity, as expected. However, it is interesting to observe that the BZI grain boundary conductivity is much improved with respect to BZY. The grain boundary conductivity is related to the grain boundary volume content and fewer boundaries will be encountered in the process of proton transport for large grainsized samples [31]. Therefore, the larger grain boundary conductivity for BZI can be attributed to the improved sinteractivity that increased grain size up to about 3 μm, effectively reducing the grain boundary volume content. Even compared with the ZnO-modified BZY (BZY-Zn) [32], the BZI sample shows better grain boundary and total conductivities (Table 1). The large BZI grain boundary conductivity, added to its good sinteractivity, makes this material interesting for BaZrO3-based fuel cell applications. The smaller conductivity than the conventional Ydoped BaZrO3 may be compensated by the better processability that may allow fabricating electrode-supported electrolyte film configuration: using tens of microns thick electrolyte films can reduce the ohmic resistance and compensate the negative effect of reduced conductivity [33]. 3.4. Chemical stability Good chemical stability is a crucial feature for SOFC electrolyte materials, especially operating with carbon-containing fuels. This issue is even more important for electrolyte films due to the reduced thickness [6,8]. Therefore, the chemical stability of BZI electrolyte layer on NiO-BZI anode half-cell was investigated. Fig. 6 shows the XRD patterns of the as-prepared BZI films and after exposure to pure
Fig. 8. Electrochemical performance (I-V and power density curves) of a BZI-based fuel cell measured from 500 to 650 °C.
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
63
Fig. 10. Typical complex-impedance plane plot of the BZI single cell measured at 600 °C at open circuit voltage.
Fig. 9. SEM cross-sectional micrograph of the BZI single cell after testing.
CO2 and boiling water. The XRD pattern of the sintered BZI (Fig. 6a) confirms that the BZI film after sintering at 1450 °C for 6 h was made of a single phase. The XRD pattern of BZI remained unchanged after CO2 exposure at 700 °C for 3 h (Fig. 6b) and after treatment in boiling water for 3 h (Fig. 6c), demonstrating a sufficient chemical stability against CO2 and against water. The chemical stability of the BZI-NiO anode was also investigated since the anode working condition can be aggressive when carboncontaining fuels are used. Moreover, stronger reactivity with CO2 and H2O may occur due to the anode porosity. Fig. 7 shows the XRD patterns of the NiO-BZI composite anode before (Fig. 7a) and after exposure to pure CO2 at 700 °C for 3 h (Fig. 7b), as well as after treatment in boiling water (Fig. 7c). Also in this case the BZI-NiO showed a good chemical stability, since no other peaks except for those of BZI and NiO could be observed. 3.5. Fuel cell performance The greatest advantage of BaZrO3-based materials for SOFC applications is their excellent chemical stability. However, the performance of BaZrO3-based fuel cells is poor, up to now. The difficult processing of Y-doped BaZrO3 allowed mostly the fabrication of electrolyte-supported fuel cells, which achieved several mW cm −2 at elevated temperatures (such as 700 °C) [34] as thick electrolyte layer (such as 1 mm) led to a large ohmic resistance for the cells, which inevitably resulted in poor cell performance. The good sinteractivity of BZI allowed fabricating anode-supported cells with reduced electrolyte thickness to tens of μm, leading to an ohmic cell resistance much smaller than that of electrolyte supported BZY cell [15,21]. Fig. 8 shows the I-V and power density curves at different temperatures for a BZI-based half-cell with a composite cathode made of PrBaCo2O5 + δ and BaZr0.7Y0.2Pr0.1O3-δ. The open circuit voltage (OCV)
was recorded at 0.985, 0.966, 0.946 and 0.929 V at 500, 550, 600 and 650 °C, respectively. These values are close to the OCV values reported in the literature for BaZrO3 electrolyte film cells [21], suggesting that the BZI electrolyte membrane was dense. With humidified hydrogen (~3% H2O) as a fuel and static air as an oxidant, the maximum power densities were 95, 84, 70 and 56 mW cm−2 at 650, 600, 550 and 500 °C, respectively. Fig. 9 shows the SEM micrograph of the single cell cross-section after testing. The BZI electrolyte layer was about 15 μm in thickness and rather dense, without any obvious pores or cracks, agreeing with the measured OCV values, as discussed above. The BZI electrolyte adhered well with both the anode and cathode layers. Table 2 compares the fuel cell performance the BZI-based cell with that of BaZrO3-based fuel cells reported in the relevant literature. Most studies for BaZrO3-based cells were performed using the electrolyte-supported configuration, and the performance of these cells was poor. Only a few studies concerning cells having BaZrO3 electrolyte films were reported and the cell performance was improved compared with that of the electrolyte-supported cells. The present BZI-based cell showed a power output of 84 mW cm −2 at 600 °C, much larger than that obtained for electrolyte-supported cells, as expected, but comparable to that of electrode-supported BZY cells. The improved sinteractivity, which enhanced proton transport at the grain boundaries and allowed sintering at lower temperatures, is the reason for the good fuel cell performance together with the acceptable conductivity of BZI. Although the current BZI cell showed larger power density than those of the most fuel cells listed in the table, a slightly larger power density output of 110 mW cm −2 at 600 °C was achieved using a Y-doped BaZrO3 4 μm-thick electrolyte film prepared by pulsed laser deposition (PLD) [40]. PLD is more suitable for smallscale applications [41], and a further reduction in the BZI electrolyte thickness may be beneficial to enhance the cell performance. Fuel cell performance is not only affected by the conductivity and thickness of the electrolyte but also by the electrode polarization, which increases with reducing the operation temperature [42]. EIS measurements at open circuit conditions of the fuel cell were used to evaluate the different resistances in the cell. Fig. 10 shows a typical EIS plot of
Table 2 Comparison of the performance of stable BaZrO3-based fuel cells reported in the literature, and the BZI-based cell developed in the present study. LSM: (La0.8Sr0.2)0.98MnO3-δ; BSCF: Ba0.5Sr0.5Co0.8Fe0.2O3-δ; LSCF:La0.6Sr0.4Co0.2Fe0.8O3-δ; PBCO:PrBaCo2O5 + δ. Year [reference]
Cell configuration
Electrolyte thickness (μm)
Cathode / Anode
Peak power density (mW cm−2)
2008 [34] 2008 [35] 2008 [15] 2009 [36] 2010 [37] 2010 [38] 2009 [36] 2009 [39] 2010 [40] Present work
Electrolyte supported Electrolyte supported Electrolyte supported Electrolyte supported Electrolyte supported Electrolyte supported Cathode supported Anode supported Anode supported Anode supported
1000 1000 600 1000 1200 300 10 30 4 15
Pt / Pt Pt / Pt Pt / Pt Pt / Pt Pt / Pt Pt / Pt LSM / NiO BSCF / NiO LSCF / NiO PBCO / NiO
7 (at 700 °C) 4 (at 700 °C) 7 (at 700 °C) 4 (at 800 °C) 3 (at 800 °C) 11 (at 800 °C) 26 (at 800 °C) 23 (at 600 °C) 110 (at 600 °C) 84 (at 600 °C)
64
L. Bi et al. / Solid State Ionics 196 (2011) 59–64
Acknowledgments This work was supported in part by the World Premier International Research Center Initiative of MEXT, Japan. The authors thank Dr. Hidehiko Tanaka for his invaluable technical assistance.
References [1] [2] [3] [4] [5] [6] [7] Fig. 11. Temperature-dependence of the BZI-based cell total, ohmic and interfacial polarization resistances, and of the ratio between ohmic resistance and cell total resistance.
the cell, measured at 600 °C. The intercept with the real axis at high frequency represents the ohmic resistance of the cell (Rohmic), which includes the electrolyte resistance and lead wires. The low frequency intercept corresponds to the total resistance of the cell. Therefore, the difference between the high frequency and low frequency intercepts with the real axis represents the total interfacial polarization resistance (Rp) of the cell. Fig. 11 shows the temperature-dependence of the total resistance, Rohmic and Rp of the cell, as determined from the EIS data. The Rohmic of the cell was 2.82, 2.39, 2.01 and 1.83 Ω cm2 at 500, 550, 600 and 650 °C, respectively, whereas the Rp values were 1.41, 0.57, 0.27 and 0.16 Ω cm2 at 500, 550, 600 and 650 °C, respectively. The ratio between Rohmic and the total resistance increased from 66% at 500 °C to 92% at 650 °C, indicating that the Rohmic makes a larger contribution to the cell total resistance in the whole testing temperature range, becoming dominant at high temperatures. Although the lower total conductivity of BZI compared with conventional Y-doped BaZrO3 partially increased Rohmic, larger electrolyte thickness and the interfacial resistances could also be the reason for the large Rohmic [11,40]. Therefore, further work, including surface modification, fabrication of electrode functional layer, and further reduction of electrolyte thickness, needs to be carried out to further improve the BZI cell performance. 4. Conclusions BaZr0.7In0.3O3-δ (BZI) was synthesized for protonic-SOFCs to overcome the poor sinteractivity of Y-doped BaZrO3, without detrimental effects on the good chemical stability. The BZI sinteractivity was improved and also the grain boundary conductivity was increased, but the bulk conductivity was lowered. Despite the total electrolyte conductivity was reduced, anode-supported fuel cell with a 15 μm thick BZI electrolyte showed a promising performance, as compared to BaZrO3-based fuel cells. The good chemical stability, improved sinteractivity and promising fuel cell performance suggested that BZI might be a good electrolyte candidate for proton-conducting SOFCs.
[8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42]
S.M. Haile, Acta Mater. 51 (2003) 5981. A. Boudghene Stambouli, E. Traversa, Renew. Sustain. Energy Rev. 6 (2002) 433. A. Bieberle-Hutter, J.L. Hertz, H.L. Tuller, Acta Mater. 56 (2008) 177. D.J.L. Brett, A. Atkinson, N.P. Brandon, S.J. Skinner, Chem. Soc. Rev. 37 (2008) 1568. H. Iwahara, H. Uchida, K. Ono, K. Ogaki, J. Electrochem. Soc. 135 (1988) 529. E. Fabbri, D. Pergolesi, E. Traversa, Chem. Soc. Rev. 39 (2010) 4355. T. Hibino, A. Hashimoto, M. Suzuki, M. Sano, J. Electrochem. Soc. 149 (2002) A1503. L. Bi, Z.T. Tao, R.R. Peng, W. Liu, J. Inorg. Mater. 25 (2010) 1. Z.M. Zhong, Solid State Ionics 178 (2007) 213. C.D. Zuo, S.W. Zha, M.L. Liu, M. Hatano, M. Uchiyama, Adv. Mater. 18 (2006) 3318. L. Bi, S.Q. Zhang, S.M. Fang, Z.T. Tao, R.R. Peng, W. Liu, Electrochem. Commun. 10 (2008) 1598. L. Bi, Z.T. Tao, C. Liu, W.P. Sun, H.Q. Wang, W. Liu, J. Membr. Sci. 336 (2009) 1. A. Tomita, K. Tsunekawa, T. Hibino, S. Teranishi, Y. Tachi, M. Sano, Solid State Ionics 177 (2006) 2951. H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, H. Suzuki, Solid State Ionics 61 (1993) 65. E. Fabbri, A. D'Epifanio, E. Di Bartolomeo, S. Licoccia, E. Traversa, Solid State Ionics 179 (2008) 558. K. Katahira, Y. Kohchi, T. Shimura, H. Iwahara, Solid State Ionics 138 (2000) 91. D. Pergolesi, E. Fabbri, A. D'Epifanio, E. Di Bartolomeo, A. Tebano, S. Sanna, S. Licoccia, G. Balestrino, E. Traversa, Nat. Mater. 9 (2010) 846. P. Babilo, S.M. Haile, J. Am. Ceram. Soc. 88 (2005) 2362. S.W. Tao, J.T.S. Irvine, Adv. Mater. 18 (2006) 1581. J.H. Tong, D. Clark, M. Hoban, R. O' Hayre, Solid State Ionics 181 (2010) 496. E. Fabbri, L. Bi, H. Tanaka, D. Pergolesi, E. Traversa, Adv. Funct. Mater. 21 (2011) 158. N. Ito, H. Matsumoto, Y. Kawasaki, S. Okada, T. Ishihara, Solid State Ionics 179 (2008) 324. S. Imashuku, T. Uda, Y. Nose, G. Taniguchi, Y. Ito, Y. Awakura, J. Electrochem. Soc. 156 (2009) B1. F. Giannici, A. Longo, A. Balerna, K.D. Kreuer, A. Martorana, Chem. Mater. 21 (2009) 2641. I. Ahmed, S.G. Eriksson, E. Ahlberg, C.S. Knee, Solid State Ionics 179 (2008) 1155. Y. Yamazaki, R. Hernandez-Sanchez, S.M. Haile, Chem. Mater. 21 (2009) 2755. T.S. Bjorheim, A. Kuwabara, I. Ahmed, R. Haugsrud, S. Stolen, T. Norby, Solid State Ionics 181 (2010) 130. V. Dusastre, J.A. Kilner, Solid State Ionics 126 (1999) 163. S.B.C. Duval, P. Holtappels, U.F. Vogt, U. Stimming, T. Graule, Fuel Cells 5 (2009) 613. E. Fabbri, D. Pergolesi, S. Licoccia, E. Traversa, Solid State Ionics 181 (2010) 1043. S.M. Haile, G. Staneff, K.H. Ryu, J. Mater. Sci. 36 (2001) 1149. S.W. Tao, J.T.S. Irvine, J. Solid State Chem. 180 (2007) 3493. N.H. Menzler, F. Tietz, S. Uhlenbruck, H.P. Buchkremer, D. Stover, J. Mater. Sci. 45 (2010) 3109. A. D'Epifanio, E. Fabbri, E. Di Bartolomeo, S. Licoccia, E. Traversa, Fuel Cells 8 (2008) 69. E. Fabbri, D. Pergolesi, A. D'Epifanio, E. Di Bartolomeo, G. Balestrino, S. Licoccia, E. Traversa, Energy Environ. Sci. 1 (2008) 355. C. Peng, J. Melnik, J.X. Li, J.L. Luo, A.R. Sanger, K.T. Chuang, J. Power Sour. 190 (2009) 447. W.M. Guo, J.A. Liu, C. Jin, J. Alloys Compd. 504 (2010) L21. A.L. Tsai, M. Kopczyk, R.J. Smith, V.H. Schmidt, Solid State Ionics 181 (2010) 1083. Y.M. Guo, Y. Lin, R. Ran, Z.P. Shao, J. Power Sour. 193 (2009) 400. D. Pergolesi, E. Fabbri, E. Traversa, Electrochem. Commun. 12 (2010) 977. E. Traversa, Interface 18 (2009) 49. E. Fabbri, D. Pergolesi, E. Traversa, Sci. Technol. Adv. Mater. 11 (2010) 044301.