Transport properties and stability of cobalt doped proton conducting oxides

Solid State Ionics 180 (2009) 160–167

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Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s e v i e r. c o m / l o c a t e / s s i

Transport properties and stability of cobalt doped proton conducting oxides Maria A. Azimova, Steven McIntosh ⁎ Department of Chemical Engineering, University of Virginia, Charlottesville, VA 22904-4741, USA

a r t i c l e

i n f o

Article history: Received 3 October 2008 Received in revised form 25 November 2008 Accepted 15 December 2008 Keywords: Proton conducting perovskite Barium cerate zirconate Transition metal cobalt doping Transport numbers Stability Intermediate temperature Solid oxide fuel cell SOFC

a b s t r a c t Cobalt doping between 2 and 10 at.% was utilized to lower the required sintering temperature of materials in the series BaCe0.5Zr0.4(Y,Yb)0.1 − yCoyO3 − δ to between 1373 and 1698 K. The required sintering temperature decreased with increasing Co content; however, significant electronic conductivity was observed in both oxidizing and reducing environments for materials with 10 at.% Co. This was accompanied by a loss of chemical stability in H2O/H2 and CO2 environments. BaCe0.5Zr0.4Yb0.07Co0.03O3 − δ was stable in these environments and provided the highest proton conductivity of the materials tested, 1.98 × 10− 3 S/cm at 923 K in humidified H2. Measurements in a hydrogen concentration cell indicated that the total ionic transference number for this material was between 0.86 and 1.00 with proton transference number between 0.84 and 0.75 at 773 and 973 K respectively. Under oxidizing conditions, the ionic transference number decreased to below 0.10. The grain boundary resistance dominated the total conductivity at low temperatures but was found to decrease with increased sintering temperature due to grain growth. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Solid oxide electrolytes find application in a number of electrochemical devices, including solid oxide fuel cells (SOFCs) and solid oxide electrolysis cells (SOECs). These devices currently require operating temperatures in excess of 973 K due to the low oxygen ion conductivity of the most common electrolyte material, yttria-stabilized zirconia (YSZ), at reduced temperatures. This high operating temperature restricts electrode material choice, accelerates thermal degradation, necessitates the use of expensive construction materials, and complicates start-up and shut-down. These issues have motivated development of intermediate temperature, 773–973 K, oxygen ion conducting electrolytes. Promising candidates such as Gdor Sm-doped CeO2 (GDC, SDC) provide high oxygen ion conductivity at reduced temperatures; however, they can also exhibit significant electronic conductivity, creating an internal ‘short circuit’ that reduces efficiency [1]. Many of the other proposed electrolytes, such as La0.9Sr0.1Ga0.8Mn0.2O3 − δ (LSGM), lack stability [2,3]. A new approach is required in order to make a significant breakthrough in lowering temperature. The focus of this paper is the solid oxide proton conductors in the series Ba(Ce,Zr)1 − x(Y,Yb)xO3 − δ; a class of materials that have been shown to provide acceptable proton conductivities in the intermediate temperature range [4–7]. Although proton conducting, these materials offer many potential advantages over both polymer electrolyte membrane fuel cells (PEMFC) (sulfur

⁎ Corresponding author. Tel.: +1 434 982 2714; fax: +1 434 982 2658. E-mail address: [email protected] (S. McIntosh). 0167-2738/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.12.013

and CO tolerance, lower cost catalysts, close thermal and system integration of the fuel reformer) and traditional SOFC (lower temperature, higher efficiency [8], simpler sealing, reduced balance of plant costs). Protons incorporated in oxides are stabilized by the electron density of nearby oxygen ions and may be considered as an effective hydroxide species, OH•O [9,10]. As such, we may consider proton incorporation as the incorporation of water into an oxygen vacancy [7,9]: H2 O + VO•• + OXO X 2OHO•

ð1Þ

Following from these arguments, proton conducting cerates and zirconates are commonly doped with trivalent cations Y3+, Yb3+, Gd3+ or Sm3+ to increase proton conductivity [7,11,12]. Eq. (1) illustrates the competition between proton incorporation and oxygen vacancy formation; dry reducing gases favor oxygen vacancy formation while humidification and high pH2 favor proton incorporation. A number of issues remain to be resolved prior to widespread application of proton conducting oxides. Acceptor doped barium cerates provide some of the highest proton conductivities [5,13–15]; however, these materials are unstable in the presence of high concentrations of water and CO2 [16–19]. Substitution of zirconium for cerium on the B-site increases the stability of the compound, but decreases the proton conductivity [17,18]. Furthermore, these materials typically require very high sintering temperatures (1873–2073 K) to form dense ceramics with high conductivity [20–22]. Finally, high resistance grain boundaries have been shown to dominate the proton conductivity at intermediate temperature [20,23]. A number of studies have suggested that the addition of small amounts of dopant cations can reduce the sintering temperature.

M.A. Azimova, S. McIntosh / Solid State Ionics 180 (2009) 160–167

Sintering temperatures have been shown to decrease to 1823 K with Gd doping and 1923 K with Nd introduction [18]. Two groups have successfully utilized ZnO addition to lower the sintering temperature to below 1723 K while maintaining high proton conductivity [24,25]. Babilo and Haile reported ion transference numbers of 0.90 at 873 K for a zinc-doped BZY material operated as an SOFC electrolyte [24]. Following this work, Tao and Irvine reported conductivity values of 0.01 S/cm at 973 K for BaCe0.5Zr0.3Y0.16Zn0.04O3 − δ [25]. In this work, we utilize Co doping between 2 and 10% of the B-site composition as method to reduce the sintering temperature of BaCe0.5Zr0.4(Y,Yb)0.1 − yCoyO3 − δ (BCZY) to below 1748 K. Co has previously been suggested as a sintering aid for CeO2 [26]. Materials were synthesized by a modified Pechini method and the resulting ceramics characterized by X-ray diffraction (XRD), scanning electron microscopy (SEM) and energy dispersive X-ray spectroscopy (EDX). The stability of the materials is investigated in humidified H2 and CO2 atmospheres at 973 K. Four-point DC conductivity was utilized to determine the total conductivity of the materials in dry and humidified 20 vol.% O2/He, He and H2 between 773 and 1173 K. The oxygen anion, proton and electronic contributions to the total conductivity of the most promising material, BaCe0.5Zr0.4Yb0.07Co0.03O3 − δ, were determined utilizing hydrogen and oxygen concentration cells. The grain and bulk contributions to the proton conductivity of this material were determined by AC impedance spectroscopy. 2. Experimental The Co doped BaCe1 − x − yZrx(Y,Yb)yO3 − δ (BCZY) materials were synthesized using a modified Pechini procedure [27]. The precursors were high purity Ba(NO3)2, Ce(NO3)2·6H2O (min 99.5%, Alfa Aesar, Ward Hill, MA, USA), Y(NO3)3·6H2O, Yb(NO3)3·5H2O, ZrO(NO3)2·6H2O (min 99% Sigma Aldrich, St. Louis, MO, USA), and Co(NO3)2·6H2O (min 99%, Acros Organics, Morris Plains, NJ, USA). Nitrate solutions were prepared by dissolving the salts in de-ionized water. Metal concentration was determined by complexometric titrations with EDTA or by thermogravimetric analysis of the water content in the nitrates [28]. These solutions were mixed in the correct ratio to form the desired composition prior to addition of EDTA, ammonium hydroxide and citric acid (min 99%, Acros Organics, Morris Plains, NJ, USA) [29]. The resulting mixture was evaporated to a gel-like state and combusted in an oven set at 573 K. Dopant levels were varied between 2 and 10% to form BaCe0.5Zr0.4Y0.08Co0.02O3 − δ(BCZY2), BaCe0.5Zr0.4Y0.07Co0.03O3 − δ (BCZY3), BaCe0.5Zr0.4Y0.05Co0.05O3 − δ (BCZY5), BaCe0.5Zr0.4Co0.10O3 − δ (BCZ10), and BaCe0.5Zr0.4Yb0.07Co0.03O3 − δ (BCZYb3). The combustion product was calcined at 1273 K for 4 h and the resulting powder ballmilled for 12 h in ethanol to achieve a homogeneous particle size. Powders were cold-pressed in a uniaxial press at a pressure of 5000– 7000 psi and sintered between 1373 and 1698 K according to Table 1. The resulting pellets were N95% of theoretical density. The heating and cooling rates were 3 K/min with 4 h dwells at the sintering temperature. Powder X-ray diffraction patterns (Scintag X-Ray Diffraction, XDS 2000, Cupertino, CA, USA) were collected on crushed pellet samples with Cu–Kα radiation source and fixed slit width. The patterns were obtained in the range of 20–65° with a 0.02° step size and a counting time of 1.4 min/degree. Rietveld structural refinements on the patterns were carried out using the GSAS package [30,31]. Powder Diffraction Files (International Centre for Diffraction Data) used to identify crystalline peaks were 00-001-0803 (BaCeO3 − δ), 01-0706758 (BaCe0.75Y0.25O3 − δ) and 01-089-2485 (BaCe0.7Zr0.3O3 − δ). To investigate the stability of the materials, powders were treated in dry and humidified H2 and CO2 atmospheres for 12 h at 973 K. All humidity levels used throughout our experiments were 3 vol.% H2O unless otherwise stated. XRD patterns were collected within 24 h of cooling to room temperature. SEM was performed using a field emission scanning electron microscope (JEOL JSM 6700, Tokyo, Japan).

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Table 1 Properties of BaCe0.5Zr0.4(Y,Yb)1 − xCoxO3 − δ materials Grain size (μm)

Composition

Short name

Sintering temperature, (K)

Lattice parameter, (Å)

% Theoretical density

BaCe0.5Zr0.4Y0.08 Co0.02O3 − δ BaCe0.5Zr0.4Y0.07 Co0.03O3 − δ BaCe0.5Zr0.4Y0.05 Co0.05O3 − δ BaCe0.5Zr0.4 Co0.10O3 − δ BaCe0.5Zr0.4Yb0.07 Co0.03O3 − δ

BCZY2

1698

4.312(3)

N95

1.5 ± 0.75

BCZY3

1648

4.309(0)

N95

1.05 ± 0.50

BCZY5

1573

4.304(4)

N95

0.62 ± 0.38

BCZ10

1373

4.302(4)

N95

0.32 ± 0.20

BCZYb3

1648

4.303(8)

N95

0.7 ± 0.32

Energy dispersive X-ray spectroscopy (EDX) was performed with Avalon system EDX attachment (PGT, Princeton, NJ). Conductivity samples were prepared by shaping sintered pellets into rectangular bars. Four Pt electrodes, two on the ends of the sample and two evenly spaced along the length of the bar, were attached using porous Pt paste (Alfa Aesar, Ward Hill, MA, USA), annealed at 1223 K in air. DC conductivity experiments were performed galvanostatically in the temperature range of 773–1173 K by a four-probe method (Reference 600 potentiostat, Gamry Instruments, Malvern, PA, USA). The maximum current was limited to 10 mA. The measurements were obtained in dry and humidified 20 vol.% O2/He, inert and H2. All gases used throughout the experiments were ultrahigh purity grade. Inert gases used were N2, He and Ar with impurity concentrations of O2 b 2 ppm and H2O b 1 ppm. The gases were humidified by passing them through a temperature controlled water bubbler. The corresponding conductivities were calculated based on the sample geometry. AC impedance spectra (Reference 600 potentiostat, Gamry Instruments, Malvern, PA, USA) were collected between 423 and 873 K by a two probe method over the frequency range 1 MHz to 0.1 Hz to measure the grain interior and grain boundary contributions to the total conductivity of these materials. Measurements were performed in humidified H2 and humidified 20 vol.% O2/He. The impedance spectra were fitted to an equivalent circuit model to determine the bulk and grain boundary resistance values using the least squares fitting procedure of EChem Analyst software (Gamry Instruments, Malvern, PA, USA). The equivalent circuit consisted of one resistor and two resistor–capacitor parallel RQ elements in series RBULK(RGBQGB) (RCPEQCPE), corresponding to bulk, grain boundary and electrode contributions, respectively [32,33]. The electrode response has a lower activation energy compared to other elements in the model and thus becomes dominant at higher temperatures. Activation energies were computed from the total conductivity data using the Arrhenius expression [24]: σ=


 A Ea exp − T RT

ð2Þ

where Ea is the activation energy, A is the pre-exponential factor and R and T have their usual meanings. Electromotive force (EMF) measurements were performed to evaluate the potential use of the material as an SOFC/SOEC electrolyte and to determine ionic and electronic transport numbers. Electrolyte pellets for these measurements were 0.5–0.7 mm thick with a diameter of 10 mm. Porous Pt (Alfa Aesar, Ward Hill, MA, USA) electrodes were attached to each side of the pellet and annealed at 1023 K in air. The resulting cathode area was 0.15 cm2. The samples were then scaled to the top of an alumina tube using Ceramabond sealant (Aremco, Valley Cottage, NY, USA). For the SOFC tests, the anode side of the cell was exposed to humidified H2 with the cathode side exposed to laboratory air. The

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Fig. 1. SEM micrograph of BCZY2 sintered at 1698 K for 4 h.

Nernst relationship was used to predict theoretical voltage for a pure proton conductor. E=

!
 0:5 pO2;cathode  pH2;anode pH2;cathode RT RT ln ln = E0 + 4F 2F pH2;anode pH2 Ocathode

peratures lower than 1698 K, as shown in Table 1. Measured XRD patterns for all of the structures were indexed to a cubic perovskite structure, space group Pm̄3m, Table 1. The cubic lattice parameter decreased with increasing Co doping levels and decreasing yttrium concentration, in agreement with previous results for yttrium-doped barium zirconates [35]. Fig. 1 is an SEM image of the surface of a typical sample, showing a well-sintered dense material. EDX was performed to probe potential dopant segregation to the grain boundaries or secondary phase formation. Multiple line scans over the polished surface of pressed pellets produced no evidence of either phenomenon, confirming the presence of the single homogeneous phase. XRD analysis was utilized to investigate the stability of these materials in the SOFC/SOEC environment. Fig. 2 shows the XRD pattern of the BCZY2 composition as-sintered (air treated) and after 12 hour exposure to humidified H2 and dry CO2 at 973 K. The material is stable in all of these atmospheres, with no additional phase observed in the post-treatment XRD patterns. All materials with Co dopant level lower than 10 at.% were found to be stable under these treatment conditions. The corresponding XRD patterns of the 10 at.% Co doped material, BCZ10, are shown in Fig. 2b. After CO2 treatment, the BCZ10 XRD pattern clearly shows an additional peak at 24.3°, associated with the formation of BaCO3 [18]. Furthermore, after exposure to humidified H2

ð3Þ

where E0 is the reversible potential, pH2, pO2 and pH2O are the partial pressures at the respective electrodes and R, T and F have their usual meanings. Proton, oxygen anion and electronic transference numbers were determined using the method of Norby et al. [34]. The EMF generated across the samples in a gas concentration cell, Eq. (4), was observed in the temperature range of 773–973 K. pH2 O I; Gas I; Ptj BCZYb3 j Pt; Gas II; pH2 O II

ð4Þ

The proton transference number was determined using O2, He and H2 and their mixtures as gases I and II in the concentration cell. By varying the partial pressure of H2 on one side of the cell while fixing it on the other and keeping the humidity constant, one can determine the total ionic transport number. This is further separated into proton, oxide ion and electron contributions by varying the humidity on one side of the cell. The gases were humidified by passing them through a temperature controlled water bubbler. The theoretical EMF of the concentration cell is given by Eqs. (5) and (6) [34]. II   RT POII 2 RT PH2 O E0 = t02− + tH + ln I −tH + ln I 4F PO 2F PH O 2 2

ð5Þ

II   RT PHII 2 RT PH2 O ln I −tO2− ln I E0 = − t02− + tH + 2F PH 2F PH O

ð6Þ

2

2

+ where t2− 0 and tH are oxygen ion and proton transference numbers respectively. The total ion transference number, ti, is given by Eq. (7).

ti = E=E0

ð7Þ

where E is the measured EMF of the concentration cell at given conditions and E0 is the theoretical EMF predicted by the Nernst relationships, Eqs. (5) and (6). Partial conductivities are resolved by multiplying the ion transference numbers by the total measured conductivity of the sample. 3. Results The modified Pechini synthesis method and Co doping resulted in dense BCZY perovskite structured materials after sintering at tem-

Fig. 2. Diffraction patterns of a) BCZY2 and b) BCZ10: i) as sintered in air, ii) after treatment in humidified CO2 and iii) after treatment in humidified H2; ⁎ denotes BaCO3 peak; ⁎⁎ denotes Ba(OH)2 peak.

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total conductivity increases linearly with the reciprocal of temperature, Fig. 3a. For example, the conductivity of BCZY3 increases from 6.0 × 10− 4 S/cm at 773 K to 1.3 × 10− 2 S/cm at 1173 K. An increasing trend in conductivity is also observed with increasing Co content. At 873 K, the total conductivities in 3 mol% humidified 20 mol% O2/ He were 7.8 × 10− 4, 1.4 × 10− 3, 3.1 × 10− 3 3.3 × 10− 2 S/cm for BCZY2, BCZY3, BCYZY5, BCZ10, respectively. The conductivity at high pO2 was not strongly influenced by substitution of Yb for Y in BCZY3 and BCZYb3. The total conductivities at 873 K of BCZY3 and BCZYb3 were 1.0 × 10− 3 and 1.2 × 10− 3 S/cm in humidified 20% O2/He, respectively. Humidification led to a decrease in total conductivity for all samples, however, the influence of humidification decreased with increasing Co content. Fig. 4, shows the influence of humidification on the conductivity of BCZY3 and BCZY5 in 20 vol.% O2/He. In both cases, the total conductivity decreases upon humidification; however, the difference is more pronounced for the 3 at.% Co doped sample. Fig. 3b shows the total conductivity in humidified inert atmosphere. This conductivity was lower in inert than in oxidizing environments for all samples. For example, the conductivity of BCZY3 in dry N2 at 973 K was 1.3 × 10− 3 S/cm compared to 4.2 × 10− 3 S/cm in 20 vol.% O2/He. In addition, the trend upon humidification was reversed when compared to the oxidizing environment, with humidification leading to an increase in total conductivity. The conductivity of BCZY3 at 973 K increased from 1.3 × 10− 3 to 2.2 × 10− 3 S/cm upon humidification of the inert atmosphere. Distinct curvature is observed with increasing temperature for the highest Co doping. The total conductivity in humidified H2 atmosphere increased with increasing Co doping, Fig. 3c. The observed conductivities were 7.8 × 10− 4, 7.7 × 10− 4, 1.0 × 10− 3 and 4.0 × 10− 3 S/cm for BCZY2, BCZY3, BCZY5 and BCZ10, respectively, at 873 K. The conductivity of all samples reach local maxima between 948 and 1023 K, beyond which they again exhibit a linear increase in conductivity to the maximum measurement temperature of 1173 K. This curvature is most pronounced for BCZ10. All DC conductivities were reproducible for the same samples except for BCZ10. As shown in the corresponding XRD pattern, Fig. 2b, this material was not stable in this severely reducing environment. Therefore, we assign the decrease in conductivity above 898 K to decomposition of the cubic perovskite structure. Substitution of Yb for Y led to a significant increase in conductivity, with a maximum

Fig. 3. Total conductivity of cobalt-doped BCZY and BCZYb3 a) in humidified 20 vol.% O2/ He, b) in humidified inert (He, N2 or Ar), and c) in humidified H2. BCZY2 (■), BCZY3 (□), BCZY5 (△), BZCY10(▲), BCZYb3 (○).

at 973 K, an additional peak at 44.5° is observed. This is associated with the formation of Ba(OH)2 [19]. Four point DC electrical conductivity measurements were conducted in the temperature range of 773 to 1173 K in humidified 20 vol.% O2/He, inert and H2, Fig. 3. In humidified 20 vol.% O2/He, the

Fig. 4. Total conductivity of Co doped BCZY in dry (BCZY3 (□), BCZY5(○)) and humidified (BCZY3 (■), BCZY5(●)) 20 vol.% O2/He.

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measured conductivity in humidified H2 of 1.9 × 10− 3 S/cm at 1023 K for BCZYb3 compared with 1.2 × 10− 3 S/cm for BCZY3. The influence of humidity upon the conductivity in reducing environments is discussed in more detail below. The total conductivity activation energies were independent of Co doping levels in humidified 20 vol.% O2/He, ranging from 0.68 eV for BCZY2 to 0.70 eV for BCZY5. These values are consistent with those previously reported for Co doped BCY compounds [36]. In humidified H2, however, the activation energies increased from 0.19 eV for BCZY2 to 0.44 eV for BCZY5, with Ea = 0.42 eV for BCZYb3, which is slightly higher than 0.35 eV value reported for pure BCZY compounds [17]. AC impedance spectra were measured for BCZYb3, the highest conductivity stable material to separate the bulk and grain contributions to the total conductivity. Fig. 5 shows representative spectra measured at 548 and 873 K. The two arcs corresponding to grain and electrode impedance are clearly differentiated at lower temperatures, Fig. 5. As the temperature is increased, the characteristic frequency of each arc increases until the grain boundary response is no longer distinguishable — above 923 K in humidified 20 vol.% O2/He. The arc corresponding to the grain interior response could not be resolved at temperatures greater than 623 K due to the characteristic frequency exceeding the maximum measurement frequency of the impedance spectrometer, 1 MHz.

Fig. 6. Grain (○) and bulk (■) conductivities of BCZYb3 sintered at 1648 K for 4 h and grain boundary (●) conductivity of BCZYb3 sintered at 1698 K for 4 h in a) humidified 20 vol.% O2/He and b) humidified H2.

Fig. 5. AC impedance arcs of BCZYb3 in humidified 20 vol.% O2/He at a) 548 K and b) 873 K.

Fig. 6, shows the grain boundary and bulk contributions to the total conductivity of BCZYb3 in humidified H2 and humidified 20 vol.% O2/ He atmospheres. The bulk conductivity is several orders of magnitude higher than the grain boundary for all temperatures and atmospheres measured. As such, the grain boundary resistance dominates the total conductivity. The corresponding bulk and grain boundary activation energies were 0.08 eV and 0.45 eV, respectively. Increasing the sintering temperature increases the measured conductivity. At 573 K in humidified H2, the grain interior conductivity of BCZYb3 sintered at 1648 K was 1.62 × 10− 2 S/cm and the grain boundary conductivity was 2.52 × 10− 5 S/cm. At the same temperature and atmosphere, the grain boundary conductivity of BCZYb3 sintered at 1698 K was 7.52 × 10− 5 S/cm. The mean grain sizes of these samples are 0.7 ± 0.32 μm and 1.5 ± 0.87 μm for lower and higher sintering temperatures, respectively. Fig. 7 shows the open circuit potential (OCP) values of the investigated materials under representative SOFC conditions of laboratory air at the cathode and humidified H2/He mixtures at the anode. All of the materials with the exception of BCZ10 provided a stable open circuit potential close to or exceeding 1.0 V at 873 K. The OCP for BCZY2 in humidified H2 was 1.01 V, decreasing with decreasing H2 concentration to 0.88 V in 3–10–87% H2O/H2/He. For BCZYb3 the OCP values in the same atmospheres decreased from 1.02 to 0.96 V. BCZ10

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Fig. 7. Open circuit potentials at 873 K of Co doped BCZY and BCZYb3 for the cell: 20 vol.% O2/He, Pt|BCZ(Y,Yb)|Pt, 3 vol.% H2O pH2, balance He. The total pressure at both electrodes was 1 atm.; BCZY2 (■), BCZY3 (□), BCZY5 (△), BCZYb3 (○). Solid line indicates the theoretical OCP predicted by the Nernst relationship (Eq. (3)).

material exhibited OCP value of 0.49 V at 873 K in humidified H2. AC impedance data obtained during these experiments was used to determine the total resistance across the cell. Accounting for the sample geometry, total conductivities were in N95% agreement with the total DC conductivity data measured in humidified H2, Fig. 3c. At 873 K, the electrolyte resistance across the BCZYb3 cell was 38.3 Ω cm2, resulting in an average conductivity of 1.57 × 10− 3 S/cm compared to 1.62 × 10− 3 S/cm measured in humidified H2 and 1.24 × 10− 3 S/cm in humidified 20 vol.% O2/He for the same material. An H2 concentration cell was utilized to determine the total ionic transference number of BCZYb3 in a hydrogen environment. One side of the cell, electrode I, was exposed to a constant atmosphere of humidified H2. The pH2 in the atmosphere at the opposite side, electrode II, was varied in a mixture of humidified H2/He. The total pressure at both electrodes was 1 atm. Fig. 8, shows the cell EMF at 773, 823 and 873 K as

Fig. 8. EMF of the hydrogen concentration cell: 3 vol.% H2O H2, Pt|BCZYb3|Pt, 3 vol.% H2O pH2, balance He. The total pressure at each electrode was 1 atm. Lines indicate the theoretical values predicted by Nernst relationship (Eq. (3)). 773 K (■ and dotted line), 823 K (□ and dashed line), 873 K (● and solid line).

165

a function of pH2 at electrode II. The EMF at all measured temperatures is close to the theoretical EMF predicted by Eq. (6). At 873 K the EMF from hydrogen concentration cell with a pH2 of 0.073 atm at electrode II was 94 mV compared to a theoretical EMF of 97 mV. This yields an ionic transference number of 0.96. Across the entire temperature range measured, the minimum ionic transference number was 0.86 at 973 K and pH2 values of 0.073 atm at electrode II and 0.776 atm at electrode I, whereas the maximum was 1.00 at 773 K with the same pH2 at electrode II and pH2 of 0.243 atm at electrode I. The total ionic transport number was studied in oxidizing atmosphere in a similar manner. The He/O2/H2O atmosphere at electrode I was fixed to provide a pO2 of 0.073 atm and pH2O of 0.03 atm. The pO2 at electrode II was varied between 0.12 and 0.97 atm in an O2/He/H2O mixture with pH2O of 0.03 atm. The total pressure at both electrodes was 1 atm and temperature was varied between 773 and 973 K. The maximum EMF value obtained was 40 mV at 773 K with a pO2 of 0.679 atm at electrode II; corresponding to a maximum ionic transference number of 0.1. At pO2 = 1 atm at electrode II and same conditions the ionic transference number decreased from 0.07 at 773 K to 0.03 at 973 K, indicating that electronic contribution dominates the transport at higher temperatures. As described in the Experimental section, separation of the total ionic transference number into proton and oxygen anion contributions was achieved using the same two-chamber system. The proton and oxygen anion transference numbers between 773 and 973 K at a pH2 of 0.073 atm at electrode II and 1 atm at electrode I are shown in Fig. 9. Across this entire range, the proton transference number was significantly higher than the oxygen anion transference number. The proton transference number decreased with increasing temperature; from a maximum of 0.84 at 773 K to a minimum of 0.75 at 973 K. The oxygen anion transport number correspondingly varied from a minimum of 0.17 at 773 K to a maximum of 0.19 at 973 K. Negligible electronic contribution of less than 0.05 was observed in this temperature range in reducing atmosphere. The electronic transference numbers were 0.01, 0.04 and 0.05 at 773, 873 and 973 K respectively. Fig. 9 shows decreasing trend of tH+ with increasing temperature. Fig. 10 shows the influence of water partial pressure on the proton transference number for a hydrogen concentration cell with H2/He/ H2O mixtures at both electrodes. The total pressure at each electrode was 1 atm, where pH2 and pH2O were 0.485 and 0.03 atm, respectively, at electrode I and 0.073 atm at electrode II with water content varying between 0.1 and 0.38 atm. The balance of the gas was He. The

Fig. 9. Proton (■) and oxygen ion (□) transport numbers measured in hydrogen concentration cell: 3 vol.% H2O H2, Pt|BCZYb3|Pt, 3 vol.% H2O pH2, balance He. The total pressure at each electrode was 1 atm.

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Fig. 10. Proton transport number of BCZYb3 at 873 K as a function of pH2O in a hydrogen concentration cell: 3 vol.% H2O pH2 = 0.485 atm, balance He, Pt|BCZYb3|Pt, pH2O, pH2 = 0.073 atm, balance He.

proton transference number increased with increasing water content from 0.79 at pH2O = 0.1 atm to 0.92 at pH2O = 0.38 atm. 4. Discussion The determined cubic perovskite structure (space group Pm̄3m) of these materials structure of is consistent with the structure of the Co free BaCe0.9 − xZrxY0.1O3 − δ parent compound [17]. The decrease in lattice parameter with increasing Co doping level is expected due to the lower ionic radius of octahedrally coordinated Co in all oxidation states compared to Y3+ [37]. The sinter temperature required to form a dense ceramic decreases with increasing Co content, in agreement with previous studies on the use of CoO as a sinter aid for CeO2 [26]. In this case, Co was added as a second phase to the parent material, and a CoO phase was reported at the grain boundaries of the sintered ceramic. Babilo and Haile [24] added ZnO to BaZr0.85Y0.15O3 − δ, and reported a similar enrichment of ZnO at the grain boundaries. This is in contrast to the work of Tao and Irvine [25] who report Zn incorporation into the lattice of BaZr0.8Y0.2O3 − δ. The primary difference between our work and these previous studies is our material synthesis approach that seeks to incorporate Co into the bulk material as a B-site dopant during wet chemical synthesis. This appears to have been successful, as our EDX scans did not find enrichment of Co at the grain boundaries of the samples. A primary issue in the development of proton conducting oxides is stability. Pure BaCeO3 is unstable in water and CO2 containing environments [19,38]; increasing Zr substitution increases stability at the expense of conductivity [7,17,18]. BaCe0.9 − xZrxY0.1O3 − δ is stable towards CO2 for x N 0.4 [17]. Our stability results indicate that increased Co doping reduces stability to both humidified H2 atmospheres and CO2. This structural instability in the humidified H2 environment may be due to the reduction of Co, and its subsequent migration out of the structure. Alternatively, instability toward H2O may be due to a higher number of oxygen vacancies available for water dissolution within the lattice with increasing Co [38]. Materials with Co content greater than 0.05 were also unstable in CO2. A second limitation to the maximum Co dopant level is the introduction of electronic charge carriers into the electrolyte material. An SOFC/SOEC electrolyte should maintain high ionic conductivity with minimal electronic conductivity in both anode (pO2 as low as 10− 24 atm) and cathode (max pO2 O(1) atm) environments. The total electrical conductivity of our materials in high pO2 (humidified and

dry 20 vol.% O2/He or He) environment increases with increasing Co doping, with BCZ10 showing the maximum conductivity at all temperatures. Furthermore, the conductivity for all of the materials increases with increased pO2. As evidenced by the low OCP for BCZ10 when utilized as a fuel cell electrolyte, the high conductivity for this material is attributable to the introduction of a significant electronic contribution to the total conductivity. The increased conductivity of both dry and humidified 20 vol.% O2/He compared to inert atmospheres indicates that p-type charge carriers dominate the transport at high pO2. Reduced conductivity upon humidification at high pO2, Fig. 4, is suggested to be due to the formation of lower mobility protons at the expense of high mobility electron holes upon humidification of a predominantly p-type conductor. The maxima in conductivity between 973 and 1073 K in the inert atmosphere may be due to reduction of Co, and associated loss of p-type charge carriers. Although the total electrical conductivity of the other materials is lower, the close to theoretical OCP values obtained suggest a primarily ionic contribution to the total conductivity under a fuel cell pO2 gradient for the lower doped materials. However, as with BCZ10, the increasing conductivity with increasing pO 2 indicates a p-type contribution at high pO2. The discrepancy between measured and theoretical OCP indicates the presence of an electronic charge carrier. As for the nature of the ionic charge carrier, the increased electrical conductivity in humidified versus dry inert and H2 atmospheres is consistent with proton transport and incorporation mechanism, as described by Eq. (1). As with high pO2 environment, we attribute the high conductivity of BCZ10 in humidified H2 to an electronic contribution, most likely n-type and due to the presence of Co2+ in this pO2 and temperature regime; however, further studies are required to confirm this. The conductivity of BCZ10 decreases sharply above 923 K, in agreement with the instability of this material observed in our XRD studies. The highest observed conductivity in humidified H2 was BCZYb3, indicating that substitution of Yb into the B-lattice is preferable over Y from the conductivity standpoint, which is consistent with previous studies performed on strontium cerates [39]. The decreases in slope and maxima in the conductivity curves at higher temperature are attributable to decreased water incorporation and, hence, proton concentration, as the temperature is increased [40]. The increasing severity of this maximum with increasing Co doping suggests an additional contribution from reduction of the Co at high temperatures. Through these initial screening studies, BCZYb3 was selected as the most promising material for more detailed study. Although a close to theoretical OCP in fuel cell environment indicates a primarily ionic conductor, it is essential to determine the nature of this ionic species and the extent of electronic conductivity in the material. The ionic transference number of BCZYb3 in H2 atmospheres is close to one; however, the material is not a pure proton conductor with proton and oxygen ion transference numbers in the range of 0.8 and 0.2. This limited oxygen ion transport may not be detrimental to fuel cell performance. Indeed, Coors demonstrated the beneficial aspect of mixed proton and oxygen anion transport by utilizing oxygen anions supplied to the anode to suppress carbon formation in a CH4 fuelled SOFC [41]. For applications in hydrogen separation or generation, however, a pure proton conductor is essential. Considering the proton incorporation mechanism, Fig. 10 shows a non-linear increase of proton transference number with increasing pH2O in a hydrogen concentration cell. This increasing trend is consistent with proton incorporation facilitated via water incorporation, Eq. (1). The non-linearity may be explained by proton formation from water and electron holes, implying that proton incorporation will be proportional to pH201/2 [42–44]. A potentially significant issue for fuel cell applications is the low ionic transference number of these materials in oxidizing atmospheres, indicating that electronic carriers, likely p-type, dominate under these conditions. Similar low but stable ionic transfer numbers

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were observed for a number of proton-conducting oxides, such as ytterbium doped strontium cerates [5] and BaCe0.9 − xZrxY0.1O3 − δ compounds [17]. High electronic conductivity in an oxidizing environment may be expected to decrease the fuel cell OCP. This is in contrast with our measured OCP and the high OCP reported in fuel cell tests [42]. The local oxygen chemical potential will set the conductivity mechanism in the electrolyte. The high OCP and match between the fuel cell electrolyte conductivity from AC impedance measurements and the conductivity of bar samples measured in reducing atmospheres suggests that the oxygen chemical potential is low throughout the electrolyte. Our measured proton transference numbers indicate that this conductivity is primarily protonic in nature in reducing atmospheres. This suggestion does, however, indicate a sharp increase in pO2 close to the exposed cathode surface. The pronounced difference observed between the grain boundary and bulk resistances indicates that the grain boundaries are the primary limitation to effective proton transport in these oxides. The difference in bulk and grain boundary activation energies also supports this conclusion. Small grain size further hinders the total conductivity of the bulk material. Grain size can be increased by sintering the materials at higher temperatures [23]. Increasing temperature by 50 K resulted in a decrease in the grain boundary resistance in both humidified 20 vol.% O2/He and H2. Observed grain boundary conductivities are lower than those of yttrium doped barium zirconate (BaZr0.9Y0.1O3 − δ), the measured values of which were about 8 × 10− 3 S/cm in humidified 20 vol.% O2/He atmosphere [35]. Finally, we compare the measured total conductivities with those reported in literature. The highest proton conductivity reported in this work for BCZYb3 is half an order of magnitude lower than that reported for pure yttrium-doped barium cerates and zirconates sintered at or above 1973 K [7,45]. Katahira et al. performed a study on various doping levels of zirconium stabilized yttrium-doped barium cerates (BaCe0.9 − x ZrxY0.1O3 − δ) [17] and reported conductivity values in humidified H2 of 4.5 × 10− 3 S/cm with pH2O = 1.7 × 103 Pa for pure BCZY after sintering at 1973 K. Based on measurements of bulk and grain boundary contributions to the conductivity, the small grain size due to low sintering temperature is the most likely source of our lower conductivity. Despite this, the conductivity of BCZYb3 was found to be comparable to that of yttria-stabilized zirconia (YSZ) below 823 K [7]. Further work is required to fabricate SOFC/SOEC from these materials. 5. Conclusions Co doping is an effective technique for reducing the required sintering temperature to form dense ceramics of proton conducting materials in the series BaCe0.5Zr0.4(Y,Yb)0.1 − yCoyO3 − δ. Increasing Co doping to 10 at.% reduced the required sinter temperature to 1373 K but introduces significant electronic conductivity and chemical instability. The materials show limited electronic conductivity and good chemical stability below 5 at.% Co. No significant Co segregation could be detected. Substitution of Yb for Y increased the proton conductivity with BaCe0.5Zr0.4Yb0.7Co0.03O3 − δ providing the highest conductivity of the materials studied. This material was a pure ion conductor with proton transference number between 0.84 and 0.75 in the temperature range

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of 773 K to 973 K in reducing humidified H2 atmospheres. At high pO2, the material shows significant p-type electronic conductivity with an ionic transference number below 0.10. Grain boundary resistance dominated the total conductivity at low temperatures. Acknowledgement Acknowledgment is made to the donors of the American Chemical Society Petroleum Research Fund for partial support of this research under grant PRF 45844-G10. References [1] [2] [3] [4] [5] [6] [7] [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] [43] [44] [45]

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