SOSI-11452; No of Pages 8
ARTICLE IN PRESS
Solid State Ionics xxx (2009) xxx–xxx
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 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
Available online xxxx
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).
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.
0167-2738/$ – see front matter © 2008 Elsevier B.V. All rights reserved.
doi:10.1016/j.ssi.2008.12.013
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013
ARTICLE IN PRESS
2
M.A. Azimova, S. McIntosh / Solid State Ionics xxx (2009) xxx–xxx
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), 01-070-6758
(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). Energy
Table 1
Properties of BaCe0.5Zr0.4(Y,Yb)1 − xCoxO3 − δ materials
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
Grain size
(μm)
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 −
RT
T
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
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013
ARTICLE IN PRESS
M.A. Azimova, S. McIntosh / Solid State Ionics xxx (2009) xxx–xxx
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=
!
pO0:5
pH2;cathode
RT
RT
2;cathode pH2;anode
= E0 +
ln
ln
pH2;anode
pH2 Ocathode
4F
2F
3
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
ln I −tH +
ln I
E0 = t02− + tH +
4F PO
2F PH O
2
2
ð5Þ
II
RT PHII 2
RT PH2 O
E0 = − t02− + tH +
ln I −tO2−
ln I
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.
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013
ARTICLE IN PRESS
4
M.A. Azimova, S. McIntosh / Solid State Ionics xxx (2009) xxx–xxx
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.
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013
ARTICLE IN PRESS
M.A. Azimova, S. McIntosh / Solid State Ionics xxx (2009) xxx–xxx
5
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
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013
ARTICLE IN PRESS
6
M.A. Azimova, S. McIntosh / Solid State Ionics xxx (2009) xxx–xxx
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).
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.
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013
ARTICLE IN PRESS
M.A. Azimova, S. McIntosh / Solid State Ionics xxx (2009) xxx–xxx
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
7
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
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013
ARTICLE IN PRESS
8
M.A. Azimova, S. McIntosh / Solid State Ionics xxx (2009) xxx–xxx
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
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]
H. Yokokawa, N. Sakai, T. Horita, K. Yamaji, M.E. Brito, MRS Bull. 30 (8) (2005) 591.
J.B. Goodenough, Ann. Rev. Mater. Res. 33 (1) (2003) 91.
S.P.S. Badwal, Solid State Ion. 143 (1) (2001) 39.
H. Iwahara, Solid State Ion. 86–88 (Part 1) (1996) 9.
H. Iwahara, T. Esaka, H. Uchida, N. Maeda, Solid State Ion. 3–4 (1981) 359.
T. Norby, Solid State Ion. 125 (1–4) (1999) 1.
K.D. Kreuer, Ann. Rev. Mater. Res. 33 (2003) 333.
A. Demin, P. Tsiakaras, Int. J. Hydrogen Energ. 26 (10) (2001) 1103.
T. Norby, M. Wideroe, R. Glockner, Y. Larring, Dalton Trans. 19 (2004) 3012.
K.D. Kreuer, Solid State Ion. 136 (2000) 149.
K.D. Kreuer, S. Adams, W. Munch, A. Fuchs, U. Klock, J. Maier, Solid State Ion. 145 (1–4)
(2001) 295.
N. Bonanos, K.S. Knight, B. Ellis, Solid State Ion. 79 (1995) 161.
K.D. Kreuer, T. Dippel, Y.M. Baikov, J. Maier, Solid State Ion. 86–8 (1996) 613.
H. Iwahara, Y. Asakura, K. Katahira, M. Tanaka, Solid State Ion. 168 (3–4) (2004)
299.
G. Ma, T. Shimura, H. Iwahara, Solid State Ion. 110 (1–2) (1998) 103.
K.D. Kreuer, Solid State Ion. 97 (1–4) (1997) 1.
K. Katahira, Y. Kohchi, T. Shimura, H. Iwahara, Solid State Ion. 138 (1–2) (2000) 91.
K.H. Ryu, S.M. Haile, Solid State Ion. 125 (1–4) (1999) 355.
S.V. Bhide, A.V. Virkar, J. Electrochem. Soc. 146 (6) (1999) 2038.
F. Iguchi, T. Yamada, N. Sata, T. Tsurui, H. Yugami, Solid State Ion. 177 (26–32)
(2006) 2381.
H. Iwahara, T. Yajima, T. Hibino, H. Ushida, J. Electrochem. Soc. 140 (6) (1993) 1687.
L. Pelletier, A. McFarlan, N. Maffei, J. Power Sources 145 (2) (2005) 262.
S.M. Haile, G. Staneff, K.H. Ryu, J. Mater. Sci. 36 (5) (2001) 1149.
P. Babilo, S.M. Haile, J. Am. Ceram. Soc. 88 (9) (2005) 2362.
S.W. Tao, J.T.S. Irvine, Adv. Mater. 18 (12) (2006) 1581.
C. Kleinlogel, L.J. Gauckler, Adv. Mater. 13 (14) (2001) 1081.
R.H.E. van Doorn, H. Kruidhof, A. Nijmeijer, L. Winnubst, A.J. Burggraaf, J. Mater.
Chem. 8 (9) (1998) 2109.
A. Vogel, Textbook of Quantitative Inorganic Analysis, Longman Scientific and
Technical, Harlow, 1986.
Z. Shao, W. Yang, Y. Cong, H. Dong, J. Tong, G. Xiong, J. Membr. Sci. 172 (1–2) (2000)
177.
H. Rietveld, J. Appl. Crystallogr. 2 (2) (1969) 65.
A.C. Larson, R.B. Von Dreele, Los Alamos National Laboratory Report LAUR, 2000, p. 86.
R.C.T. Slade, S.D. Flint, N. Singh, Solid State Ion. 82 (3–4) (1995) 135.
K.D. Kreuer, Solid State Ion. 125 (1–4) (1999) 285.
D.P. Sutija, T. Norby, P. Björnbom, Solid State Ion. 77 (1995) 167.
F. Iguchi, N. Sata, T. Tsurui, H. Yugami, Solid State Ion. 178 (7–10) (2007) 691.
T. Shimura, H. Tanaka, H. Matsumoto, T. Yogo, Solid State Ion. 176 (39–40) (2005)
2945.
R.D. Shannon, Acta Crystallogr. A32 (1976) 751.
C.W. Tanner, A.V. Virkar, J. Electrochem. Soc. 143 (4) (1996) 1386.
T. Yajima, H. Suzuki, T. Yogo, H. Iwahara, Solid State Ion. 51 (1–2) (1992) 101.
T. Norby, N. Christiansen, Solid State Ion. 77 (1995) 240.
W.G. Coors, J. Power Sources 118 (1–2) (2003) 150.
H. Iwahara, H. Uchida, S. Tanaka, Solid State Ion. 9–10 (Part 2) (1983) 1021.
S. Stotz, C. Wagner, Berich. Bunsen Gesell. 70 (1966) 781.
D.A. Shores, R.A. Rapp, J. Electrochem. Soc. 119 (1972) 300.
N. Bonanos, B. Ellis, K.S. Knight, M.N. Mahmood, Solid State Ion. 35 (1–2) (1989)
179.
Please cite this article as: M.A. Azimova, S. McIntosh, Solid State Ionics (2009), doi:10.1016/j.ssi.2008.12.013