Embed Size (px)
Citation preview
www.elsevier.com/locate/ssi
Solid State Ionics 176 (
Ferrite-based perovskites as cathode materials for anode-supported
solid oxide fuel cellsB
Part I. Variation of composition
Andreas Mai, Vincent A.C. Haanappel, Sven Uhlenbruck, Frank Tietz*, Detlev Stover
Institute for Materials and Processes in Energy Systems, Forschungszentrum Julich, IWV-1, D-52425 Julich, Germany
Received 16 August 2004; received in revised form 9 March 2005; accepted 10 March 2005
Abstract
The properties and the applicability of iron- and cobalt-containing perovskites were evaluated as cathodes for solid oxide fuel cells
(SOFCs) in comparison to state-of-the-art manganite-based perovskites. The materials examined were La1�x�ySrxCo0.2Fe0.8O3� d (x =0.2
and 0.4; y =0�0.05), La0.8Sr0.2FeO3� d, La0.7Ba0.3Co0.2Fe0.8O3� d and Ce0.05Sr0.95Co0.2Fe0.8O3� d. The main emphasis was placed on the
electrochemical properties of the materials, which were investigated on planar anode-supported SOFCs with 8 mol% yttria-stabilised zirconia
(8YSZ) electrolytes. An interlayer of the composition Ce0.8Gd0.2O2� d was placed between the electrolyte and the cathode to prevent
undesired chemical reactions between the materials. The sintering temperatures of the cathodes were adapted for each of the materials to
obtain similar microstructures. In comparison to the SOFCs with state-of-the-art manganite-based cathodes, SOFCs with
La1� x� ySrxCo0.2Fe0.8O3� d cathodes achieved much higher current densities. Small A-site deficiency and high strontium content had a
particularly positive effect on cell performance. The measured current densities of cells with these cathodes were as high as 1.76 A/cm2 at
800 -C and 0.7 V, which is about twice the current density of cells with LSM/YSZ cathodes. SOFCs with La0.58Sr0.4Co0.2Fe0.8O3� d cathodes
have been operated for more than 5000 h in endurance tests with a degradation of 1.0–1.5% per 1000 h.
D 2005 Elsevier B.V. All rights reserved.
PACS: 84.60.Dn (Electrochemical conversion and storage); 81.05.Je (Ceramics and refractories)
Keywords: Solid oxide fuel cells; SOFC; Cathode; Perovskite; LSCF; (La,Sr)(Co,Fe)O3; Lanthanum ferrite
1. Introduction
Reducing the costs for solid oxide fuel cell (SOFC)
systems is one of the main current issues of this technology
[1,2]. Therefore, efforts are being made to lower the
operating temperature below the commonly used 800–
1000 -C and to increase the power density of the SOFCs.
Operating temperatures of 750 -C or below make it possible
to use cheaper interconnect materials and a wider range of
sealing materials [4]. It also reduces the degradation of stack
0167-2738/$ - see front matter D 2005 Elsevier B.V. All rights reserved.
doi:10.1016/j.ssi.2005.03.009
i In memoriam Professor B.C.H. Steele.
* Corresponding author. Tel.: +49 2461 61 5007; fax: +49 2461 61 2455.
E-mail address: [email protected] (F. Tietz).
materials and therefore leads to improved reliability and
long-term stability.
The decrease of the power densities at lower temper-
atures is mainly due to overpotentials at the cathode [3,5,6].
This makes it necessary to develop new cathode materials
with a higher electrocatalytic activity than those of the state-
of-the-art La1� x � ySrxMnO3� d (LSM) perovskites. In
addition to a high electronic conductivity, these materials
should have high oxygen ion conductivity to enlarge the
area where oxygen reduction can take place as well as high
oxygen surface exchange coefficients for faster kinetics at
the gas/cathode interface. Furthermore, no chemical reac-
tions with the surrounding materials should occur, and the
thermal expansion coefficient (TEC) should be close to that
of the electrolyte to avoid mechanical stresses.
2005) 1341 – 1350
A. Mai et al. / Solid State Ionics 176 (2005) 1341–13501342
A group of materials that fulfil some of these require-
ments are iron- and cobalt-containing perovskites, for
example La1� xSrxCo1� yFeyO3� d (LSCF), that have
already been known for their high oxygen permeability [7]
and high electrocatalytic activity (see, e.g., Refs. [8–10]).
However, these materials have to be selected carefully
because they have a significantly higher thermal expansion
coefficient than the commonly used 8 mol% yttria-stabilised
(8YSZ) electrolyte [11,12]. Furthermore, these type of
perovskites form SrZrO3 with the 8YSZ electrolyte at high
temperatures [10,11]. To overcome these problems, an
interlayer consisting of Ce0.8Gd0.2O2� d (CGO) between
cathode and electrolyte can be used [13,14].
In this first part on ferrite-based cathodes materials, the
applicability of several mixed-conducting perovskites is
evaluated as cathode materials for SOFCs with an 8YSZ
electrolyte. Their properties are compared with the state-of-
the-art LSM/YSZ composite cathodes. The processing of
the CGO interlayer and its influence on performance will be
the subject of a forthcoming paper, the second part in this
series.
2. Experimental
2.1. Powder preparation
Most of the cathode materials used were synthesised by
spray-drying as described by Kontouros et al. [15]. One of
the powders (La0.8Sr0.2Co0.2Fe0.8O3� d) was also synthes-
ised by a citrate complexation (Pechini) route [16], while
La0.7Ba0.3Co0.2Fe0.8O3� d was only synthesised by citrate
complexation. In addition to these powders, La0.58Sr0.4Co0.2Fe0.8O3� d was also provided by ECN (Energy
Research Centre of the Netherlands, Petten) while
Ce0.05Sr0.95Co0.2Fe0.8O3� d was provided by the Dresden
University of Technology (Germany). All cathode powders
were calcined at 700–900 -C in order to develop the
perovskite phase. The abbreviations used hereafter for the
different compositions are listed in Table 1.
Table 1
Data of cathode materials used in this paper
Abbreviation Composition TEC (�10�6 K�1)
(30–1000 -C)
LSM La0.65Sr0.3MnO3� d 12.3 [19]
L55SCF La0.55Sr0.4Co0.2Fe0.8O3� d 17.1
L58SCF La0.58Sr0.4Co0.2Fe0.8O3� d 17.4
L60SCF La0.6Sr0.4Co0.2Fe0.8O3� d 17.5 [12]
L78SCF La0.78Sr0.2Co0.2Fe0.8O3� d 13.8
L80SCF La0.8Sr0.2Co0.2Fe0.8O3� d 14.8 [12]
L70BCF La0.7Ba0.3Co0.2Fe0.8O3� d 16.8
CSCF Ce0.05Sr0.95Co0.2Fe0.8O3� d 23.7
L80SF La0.8Sr0.2FeO3� d 11.9
* P=perovskite, (S)= traces of spinel, (L)= traces of La2O3.
The stoichiometry of the powders was controlled by
optical emission spectroscopy (ICP-OES), while the phase
composition was evaluated by X-ray diffraction using a
Siemens D500 equipped with a monochromated Cu Ka
radiation source. The XRD patterns showed traces of La2O3,
La(OH)3 and of a spinel-type formation (Co,Fe)3O4 for
some of the powders that were calcined at 900 -C or below.
After calcination at 1100 -C, traces of La2O3 were only
found in L80SF, while all the other powders did not show
any La2O3 or La(OH)3. For the La0.55Sr0.4Co0.2Fe0.8O3� d
with 5% A-site deficiency, traces of the spinel were visible
in the XRD pattern after calcination at 1100 -C. These tracesof (Co,Fe)3O4 were not considered to be detrimental for use
as an SOFC cathode. After calcination of the powders, they
were ground by ball milling for several hours until a mean
particle size (d50) of approx. 0.8 Am was achieved.
Ce0.8Gd0.2O2� d (CGO) was obtained from Treibacher
Auermet (Austria) and was also ground by ball milling.
2.2. Sample preparation
All the cathode materials were tested on anode-supported
single cells consisting of an anode substrate (Ni/8YSZ) with
an average thickness of about 1.5 mm. The substrates with a
size of 50�50 mm2 were produced by warm pressing using
a so-called Coat-Mix material [17]. An electrochemically
active anode functional layer (Ni/8YSZ, thickness approx.
10 Am) and an electrolyte (8YSZ, thickness approx. 10 Am)
were both deposited by vacuum slip casting on de-bindered
substrates and co-fired at 1400 -C. More details about the
manufacturing process can be found elsewhere [2,18].
The subsequent layers were screen-printed using pastes
based on the above-mentioned ceramic powders, an ethyl
cellulose binder and a terpineol-based solvent. The area
of the cathode layers was 40�40 mm2. The cells used
for a comparison to the reference cathode material
La0.65Sr0.3MnO3� d (LSM) had a LSM/YSZ cathode func-
tional layer and a LSM current collector as described in Ref.
[18]. For all cells with ferrite-based cathodes, a CGO
interlayer preventing chemical reactions between the cath-
Sintering temperature Crystalline phases* formed afte
calcination at 1100 -C
1100 -C P
1080 -C P, (S)
1080 -C P
1200 -C P
1060 -C P
Spray-dried: 1150 -C P
Pechini: 1080 -C P
1120 -C P
1000 -C P
1150 -C P, (L)
r
A. Mai et al. / Solid State Ionics 176 (2005) 1341–1350 1343
ode and the 8YSZ electrolyte was applied on the electrolyte
and sintered at 1300 -C. Some of the cells were prepared
with an interlayer having a thickness of 5 Am and consisting
of coarser CGO powder (d50=0.89 Am) (hereafter denoted
as interlayer type I), while the others had a thickness of 7
Am and finer CGO powder (d50=0.2 Am) (interlayer type
II). The cathode was then printed and sintered on the
interlayer resulting in a thickness of 45 Am. For the sintering
temperatures of the cathodes, see Table 1 and Section 3.2.
2.3. Characterisation of the structural and chemical
properties
Characterisation of the microstructures was performed
using scanning electron microscopes (Jeol T300 and LEO
1530 (Gemini)) equipped with energy dispersive X-ray
(EDX) analysis systems.
Thermal expansion measurements on sintered dense
bodies of the cathode materials were carried out as described
in Ref. [19].
The element distribution was examined by EDX and
SIMS (secondary ion mass spectroscopy, ION-TOF Tof-
SIMS) on polished cross sections of the samples.
2.4. Electrochemical characterisation
Electrochemical measurements of single cells were
performed in an alumina test housing placed inside a
furnace. In order to obtain sufficient electronic contact
between the cell and the electronic devices a Ni mesh and a
Pt mesh were used at the anode side and the cathode side,
respectively. Air tightness of the gas compartment was
obtained by a gold sealant. During the start-up of the tests an
argon flow was injected at the anode side and an air flow at
the cathode side. The temperature was then slowly increased
to the temperature for anode reduction. After reaching this
temperature, the anodes of the single cells were reduced by
gradually replacing the argon with hydrogen. Water vapour
(3 vol.%) was added by saturating the hydrogen gas through
a water bubbler and condenser (supersaturation and con-
densation) at the desired dew point of 24 -C. The total gas
flows of hydrogen and air were both set at 1000 ml/min
(standard temperature and pressure: STP) using mass flow
Fig. 1. L58SCF cathodes sintered at 1040 (left), 1080 (middle) and 1120 -C (righ
type I and the electrolyte YSZ.
controllers. The electrochemical performance was measured
between 650 and 800 -C. All electrochemical data were
obtained by DC methods using a current-control power
supply type Gossen 62N-SSP500-40 (Gossen-Metrawatt
GmbH, Germany), and a computer-controlled data acquis-
ition system including a data logger type NetDAQ 2640A
(Fluke, The Netherlands). The current–voltage character-
istics were measured with increasing current load by a
sequential step change of 0.0625 A/cm2 starting from zero
until either the voltage dropped below 0.7 V or the
maximum current load of 1.25 A/cm2 was reached. A
comparison of the electrochemical performance regarding
the different types of single cells was made by comparing
the current densities at 0.7 V. Calculations of the current
density at 700 mV and 650, 700, 750, 800 -C are based on
inter- or extrapolation. Calculations of the area-specific
resistance are based on linear regression of the current–
voltage curves at 0.7 V. For all different types of single cells,
at least two nominally identical cells were measured. The
data given in the tables is the average value of the
measurements, while the error margins show the differences
between these identically prepared cells.
3. Results and discussion
3.1. Thermal expansion coefficients
Iron- and cobalt-based perovskites are known to have
thermal expansion coefficients (TEC) that are well above
the TEC of the 8YSZ electrolyte (aYSZ=10.8�10�6 K�1
(30–1000 -C), [11]). The TECs of the materials used in the
present work as measured by dilatometry are shown in Table
1. For most of the powders, the slope of the TEC curves
increases with temperature, which can be explained by a
loss of oxygen at elevated temperatures [20,21]. The A-site
deficiency has only a small effect on the TEC, while a
higher Sr content results in higher TECs due to higher
oxygen vacancy concentrations, as has already been out-
lined in Ref. [21].
The cathodes made of CSCF and L60SCF showed cracks
and were partially spalled off from the CGO interlayer when
sintered on the 50�50 mm2 cells. This was considered to
t). All micrographs show the cathode on top followed by a CGO interlayer
Fig. 3. Microstructure of an LSM/YSZ composite cathode on a YSZ
electrolyte.
A. Mai et al. / Solid State Ionics 176 (2005) 1341–13501344
result from thermal stresses due to the high TEC of the
materials, which was aggravated by the high sintering
temperature in the case of the L60SCF. For the other types
of cathodes only microcracks were visible on the surface.
For these cathodes, the porous CGO interlayer with a
slightly higher TEC (12.8�10�6 K�1 (20–1000 -C), [19])than 8YSZ might reduce the mechanical stresses on the
cathode preventing spallation of the cathode layer.
3.2. Sintering behaviour of the different materials
It has already been shown on LSM cathodes that the
microstructure has a strong influence on the performance
of the cathode [5] and that it can be tuned by varying the
sintering temperature. A coarser structure improves factors
like ionic and electronic conductivity and gas permeability
of the cathode, while a finer structure leads to a higher
specific surface area of the cathode and therefore to a
greater number of reaction sites. Often, when comparing
different cathode materials, the sintering temperature was
optimised beforehand for one of the materials under
investigation [10,22,23] and all of the various composi-
tions were subsequently sintered at that temperature. As a
consequence, the materials’ influence on the electrochem-
ical performance of the SOFC could not be easily
identified.
Fig. 1 shows three cathode layers consisting of L58SCF
sintered at 1040, 1080 and 1120 -C, for 3 h each. The three
pictures show a coarsening of the microstructure with
increasing temperature due to enhanced sintering of the
powder particles.
The sintering activity of the materials is, however,
already very different when the stoichiometry is slightly
changed. In Fig. 2, the effect of the A-site deficiency of the
cathode material on the sintering activity is shown. While
the L60SCF powder has hardly formed any sintering necks
at 1080 -C, the L55SCF particles are in good contact to each
other, the latter appearing similar to those of the L58SCF
cathode in Fig. 1b. Sintering the L60SFC cathode at 1200
-C leads to a microstructure that is similar to the above-
mentioned L58SFC and L55SCF cathodes sintered at 1080
-C as is shown in Fig. 2c. In Fig. 3, the microstructure of an
LSM/YSZ composite cathode is shown for comparison.
Fig. 2. Cathodes consisting of L55SCF and L60SCF sintered at 1080 -C (left and m
micrographs show the cathode on top followed by a CGO interlayer type I and t
Unless noted otherwise, an appropriate sintering temper-
ature was chosen so that the cathode of all tested cells had a
microstructure similar to the L58SCF cathode sintered at
1080 -C (Fig. 1b). The sintering temperatures for all
cathodes investigated here are listed in Table 1. For the
two L80SCF powders which were differently synthesised
(citrate complexation and spray-drying), the temperature
was also adjusted to the different sintering activities.
3.3. Chemical reaction with 8YSZ
As has already been previously reported, LSCF reacts
with 8YSZ at high temperatures forming the insulating
SrZrO3 [10,24]. This can be partially prevented by an
interlayer of an additional material, for example CGO,
between cathode and electrolyte. Due to the low sintering
activity of CGO when compared with 8YSZ, the interlayer
is porous. This allows the Sr ions to diffuse through the
interlayer and form an insulating SrZrO3 barrier between
CGO and 8YSZ. The formation of this barrier is increased
by two parameters: high sintering temperatures of the
cathode and high strontium content of the cathode material.
In general, higher Sr content leads to a less stable perovskite
[25], resulting in an increased SrZrO3 formation. As can be
seen in the EDX mappings in Fig. 4, the cell with L80SCF
cathode showed no detectable SrZrO3 formation, while for
the L58SCF cell, a small zone of a Sr-rich phase can be
detected close to the electrolyte. The slightly elevated Sr-
signal in the YSZ-electrolyte is an artefact due to an overlap
of the La-series in the characteristic X-ray spectra of
iddle, respectively) and an L60SCF cathode sintered at 1200 -C (right). All
he electrolyte YSZ.
0.0 0.2 0.4 0.6 0.8 1.0 1.20.6
0.7
0.8
0.9
1.0
1.1
750 °C
Cel
l vol
tage
(V
)
Current density (A/cm2)
TS= 1040 °C
TS= 1080 °C
TS= 1120 °C
Fig. 5. Current–voltage curves at 750 -C of single cells with L58SCF-
based cathodes sintered at various temperatures including a CGO interlayer
type I.
A. Mai et al. / Solid State Ionics 176 (2005) 1341–1350 1345
strontium and zirconium. Because of this superposition, the
polished cross sections were additionally examined by
SIMS. In accordance with the EDX mappings, there was
no SrZrO3 detectable for the perovskites with 20% Sr on the
A-site (L80SCF, L78SCF, L80SF), while there was a small
zone of the Sr-rich phase detectable for all of the materials
with 40% Sr on the A-site (L55SCF, L58SCF, L60SCF).
The A-site deficiency did not result in any definite effect on
SrZrO3 formation. This is in contrast to findings on
manganite-based perovskites [26], where an A-site defi-
ciency led to a higher stability with respect to strontium
depletion. There was also an inhomogeneous distribution of
cobalt visible for the L58SCF and L55SCF cathodes. This
was considered to result from cobalt oxide in the cathode (as
in Ref. [27]) and was not thought to be harmful to the
cathode’s performance.
3.4. Electrochemical performance of the materials
3.4.1. Influence of sintering temperature
As already mentioned, the sintering temperature has a
strong influence on the microstructure of the cathodes and
therefore influences the electrochemical performance of the
cells [5]. Consequently, the electrochemical performance of
cells with an L58SCF cathode sintered at 1040, 1080 and
1120 -C was tested to find the optimum in electrochemical
performance. These cells contain a CGO interlayer of type I
(mean particle size 0.89 Am, thickness 5 Am). As can be seen
in Fig. 5 and in Table 2, the electrochemical performance is
significantly lower after sintering at 1120 -C than after 1080
-C, which can be explained by the smaller intrinsic surface
area of the cathode due to particle growth. The higher
sintering temperature also increases the above-mentioned
SrZrO3 formation, which lowers the electrochemical per-
formance of the cathodes. A lower sintering temperature
(1040 -C) leads to a slightly lower current density at 800 -Cand a slightly higher one at 650 -C. This can be explained bya poorer adhesion between cathode and interlayer or among
the cathode particles themselves, which compensates the
effect of a higher surface area for the samples sintered at a
lower temperature. However, the difference between these
two is too small to draw final conclusions.
Fig. 4. Formation of a Sr-rich phase at the electrolyte: EDX line scans for stronti
electrolytes, CGO interlayers, and (a) L80SCF, (b) L58SCF cathode. The baselin
3.4.2. Influence of A-site deficiency and synthesis route
A-site deficiency influences the sintering activity and
also the stability with regard to Sr depletion. As mentioned
above, the sintering temperature was adapted to the different
sintering activities to exclude an influence of the micro-
structure. In Table 3 and Fig. 6, the results of the
electrochemical measurements of cells with L55SCF,
L58SCF, L60SCF (all with x =0.4), and L78SCF and
L80SCF (with x =0.2) cathodes are shown. The difference
of about 10% between the values for the L58SCF cells listed
in Tables 2 and 3 is due to the different CGO layers (see
Section 2.2).
The performance is significantly higher for the A-site
deficient than for the non-deficient materials. The main
reason for this might be that the oxygen defect concentration
is higher for the A-site deficient than for the non-deficient
powders [28]. For L60SCF, the cell performance also
decreased due to partial spallation of the cathode during
the measurement, which is thought to result from higher
mechanical stresses due to the sintering temperature of 1200
-C. This is in contrast to Simner et al. [23], who found that
La0.8Sr0.2FeO3� d has a higher performance than their A-site
or B-site deficient equivalents. However, here the same
sintering temperature was used for all materials, so that the
um, cerium and zirconium taken along the black lines of SOFCs with YSZ
e of the Zr and Ce lines are shifted for better readability.
Table 2
Current density (A/cm2, at 700 mV) and area-specific resistance (mV cm2) of L58SCF-based single cells as a function of the sintering temperature TS of the
cathode
Temperature (-C) Current density (A/cm2, 700 mV) Area-specific resistance (mV cm2)
TS=1040 -C 1080 -C 1120 -C TS=1040 -C 1080 -C 1120 -C
800 1.51T0.12 1.60T0.08 1.00T0.10 215T2 195T3 307T2
750 1.26T0.10 1.31T0.04 0.79T0.09 272T3 239T3 409T4700 0.93T0.04 0.92T0.01 0.52T0.05 360T7 326T4 638T7
650 0.60T0.01 0.55T0.01 0.29T0.01 511T7 531T16 1275T37
A. Mai et al. / Solid State Ionics 176 (2005) 1341–13501346
cells with deficient cation stoichiometry were probably
over-sintered or poorly adhered.
Performance of the materials with the same A-site
deficiency was generally higher with higher Sr content. A
higher amount of Sr atoms instead of the trivalent lanthanum
on the A-site is known to increase the ionic and electronic
conductivity and the surface exchange of oxygen, which can
be explained by the larger number of oxygen vacancies and
electronic holes [21,29,30]. However, the thermal expansion
coefficient also increases with higher Sr content [12,21],
which can cause mechanical problems. The electrochemical
data in Table 3 and Fig. 6 show that the materials with 40%
Sr on the A-site (L55SCF, L58SCF, L60SCF) perform much
better than the corresponding perovskites with 20% Sr
substitution (L78SCF, L80SCF). Of the compositions
investigated, the cells with a L58SCF cathode achieved the
highest current densities of 1.76 A/cm2, at 800 -C and 0.7 V.
For L80SCF, cathodes made of powders synthesised by
two routes were tested. In addition to the spray-dried
powder, another powder was synthesised by the Pechini
route. The citrate complexation (Pechini) method leads to a
powder with very fine particles and high specific surface
area. Also, a calcination temperature of 700 -C was
sufficient to produce a phase-pure powder. This resulted in
a higher sintering activity, and therefore a lower sintering
temperature than for the spray-dried powder had to be
chosen (1080 instead of 1150 -C). When comparing the
Table 3
Current density (A/cm2, at 700 mV) and area-specific resistance (mV cm2) of LSCF
route
Temperature (-C) Synthesis by spray-drying
CGO layer type II
L55SCF L58SCF L60SCF L
Current density (A/cm2, 700 mV)
800 1.23T0.04 1.76T0.08 0.90T0.04 1
750 1.04T0.04 1.43T0.08 0.77T0.02 0
700 0.74T0.03 0.99T0.06 0.54T0.02 0
650 0.46T0.02 0.58T0.03 0.34T0.01 0
Area-specific resistance (mX cm2)
800 253T5 179T2 325T3 3
750 302T2 219T3 437T3 4
700 431T9 297T4 650T6 5
650 694T12 522T10 1076T29 9
+ Results from only one cell.
electrochemical performance, however, the difference of the
two materials is negligible. Nevertheless, the lower sintering
temperature for the Pechini powder could lower SrZrO3
formation and therefore the Pechini method might be an
option for other compositions, e.g. L58SCF. In addition to
the L58SCF synthesised by spray drying, we tested L58SCF
provided by ECN (Petten, Netherlands). The performance of
cells with cathodes made of this L58SCF was lower than
using the spray-dried L58SCF powder (Table 3).
3.4.3. Influence of perovskite composition
Additionally to the LSCF materials, other perovskites
were tested such as L70BCF with barium instead of
strontium, L80SF without cobalt on the B-site, CSCF with
strontium and cerium on the A-site as well as the state-of-
the-art material LSM. For all of the cells except the LSM-
based cells, a CGO interlayer type I was used.
For L70BCF a high contact resistance was measured, and
therefore an additional LSM current collector was applied
on top of the L70BCF layer. With respect to the electro-
chemical data (Table 4, Fig. 7), its performance is
approximately the same as the corresponding L80SCF
material. Thus, substitution of the alkaline atom seems to
have little influence on performance. Together with the
findings of Ahmad-Khanlou [31] and a more recent paper
by Qiu et al. [32] reporting that replacing lanthanum by
other lanthanides has little influence on electrochemical
-based single cells with different A-site deficiency, Sr content and synthesis
CGO layer type I
78SCF+ L80SCF L58SCF powder
from ECN
L80SCF citrate
complexation
.20 0.68T0.01 1.28T0.07 0.74T0.04
.96 0.56T0.01 1.10T0.03 0.59T0.04
.63 0.42T0.01 0.75T0.01 0.41T0.03
.36 0.24T0.01 0.49T0.01 0.25T0.01
32 493T4 239T2 399T6
09 546T4 283T3 540T9
57 744T16 444T7 777T10
76 1228T37 680T12 1251T17
Fig. 6. Current–voltage curves at 750 -C of single cells with LSCF cathodes containing 40% (left) and 20% Sr (right) on the A-site and different levels of
deficiency. The cathode powders were synthesised by spray-drying except L80SCF-CC, which was synthesised by citrate complexation.
A. Mai et al. / Solid State Ionics 176 (2005) 1341–1350 1347
behaviour, this points to a low influence of the A-site atoms
in general.
The type of B-site atom is generally known to have a
stronger influence on the electrochemical behaviour.
Therefore, L80SF was tested in comparison to L80SCF,
with the latter having 20% cobalt on the B-site instead of
iron. With L80SF lower current densities were achieved
over the full range of temperatures. This can be explained
by the beneficial effects of cobalt for the catalytic activity
for oxygen reduction resulting in a higher performance of
L80SCF. In addition to the lanthanum-based perovskites
the highly oxygen-deficient CSCF was tested. The oxygen
ion mobility and surface exchange coefficients for these
kinds of materials are much higher than for the lanthanum-
based materials and they were proposed as a high-
performance cathode material by Trofimenko and Ullmann
[33], and by Colomer et al. [34]. The drawbacks of these
materials are, however, quite obvious when using them as a
cathode in an SOFC: although a rather low sintering
temperature could be chosen (1000 -C) to achieve the
desired microstructure, the mechanical stresses due to the
high TEC mismatch lead to partial detachment of the
cathode before and after the electrochemical character-
isation. As a consequence, the electrochemical performance
Table 4
Current density (A/cm2, at 700 mV) and area-specific resistance (mV cm2)
of single cells with different perovskite-based cathode materials
Temperature (-C) L70BCF L80SF CSCF LSM/YSZ
Current density (A/cm2, 700 mV)
800 0.65T0.05 0.51T0.08 0.64T0.11 0.92T0.05750 0.48T0.03 0.47T0.07 0.49T0.09 0.55T0.11
700 0.30T0.01 0.36T0.05 0.34T0.06 0.36T0.08
650 0.17T0.01 0.24T0.04 0.20T0.03 0.22T0.03
Area-specific resistance (mX cm2)
800 476T3 586T6 520T3 343T64
750 590T18 714T11 644T8 517T174
700 821T19 936T7 938T28 767T248650 1444T44 1408T19 1704T74 1327T312
is much lower (see Table 4 and Fig. 7) than one would
expect from the properties of this material.
For reasons of comparison, the performance of a state-of-
the-art LSM/YSZ composite cathode, which is used as the
standard material for stack tests at Forschungszentrum
Julich, is shown in Fig. 7. The current densities of the
better performing LSCFs (L55SCF, L58SCF, L78SCF) are
up to two times higher than for the LSM cathode, while the
performance of the other perovskites is about the same or
lower than that of LSM.
3.5. Behaviour of L58SCF cathodes under various
measurement conditions
Before using a new cathode material (i.e. replacing
LSM by LSCF) in SOFC stacks, it has to be ensured that
the material does not fail under stack operating conditions.
Single cells with a cathode made of L58SCF from ECN
and a CGO interlayer type II were therefore tested under
various experimental conditions: long-term performance,
with methane and various water vapour contents in the
fuel gas.
0.0 0.2 0.4 0.6 0.8 1.0 1.20.6
0.7
0.8
0.9
1.0
1.1 L80SCF L80SF L70BCF CSCF LSM/YSZ
750 °C
Cel
l vol
tage
(V
)
Current density (A/cm2)
Fig. 7. Current–voltage curves of single cells at 750 -C with different
perovskite compositions used as cathode material.
0.0 0.2 0.4 0.6 0.8 1.0 1.20.6
0.7
0.8
0.9
1.0
1.1 3 vol.% H2O
25 vol.% H2O
50 vol.% H2O
Cel
l vol
tage
(V
)
Current density (A/cm2)
750 °C
Fig. 8. Current–voltage curves of an SOFC with an L58SCF-based cathode
as a function of the water vapour concentration (fuel gas: H2=1000 ml/min,
oxidant: air=1000 ml/min).
0.0 0.2 0.4 0.6 0.8 1.0 1.20.6
0.7
0.8
0.9
1.0
1.1 33 vol.% CH4 + 67 vol.% H
2O
Cel
l vol
tage
(V
)
Current density (A/cm2)
797 °C 757 °C 698 °C 643 °C
Fig. 9. Current–voltage curves of an SOFC with an L58SCF-based cathode
as a function of temperature with methane as fuel gas: CH4=330 ml/min,
H2O=670 ml/min; oxidant: air=1000 ml/min.
A. Mai et al. / Solid State Ionics 176 (2005) 1341–13501348
3.5.1. Effect of H2O concentration with constant hydrogen
flow
The effect of H2O concentration on the cell voltage and
power output at 750 -C with a constant hydrogen flow
(1000 ml/min) is shown in Fig. 8. With increasing water
concentration the OCV decreased, which is obvious due to a
lower Nernst potential, as the oxygen partial pressure
increases with higher H2O content in the fuel gas. The
average current density also slightly decreased with higher
amounts of water (see Table 5). The area-specific resistance
did not significantly change with different water concen-
trations. A similar behaviour of the electrochemical data was
obtained with other test temperatures.
No obvious differences were found with different H2O
concentrations including a constant total gas flow of 1000
ml/min (STP).
3.5.2. Effect of the presence of methane
In addition to H2–H2O gas mixtures, experiments were
performed with methane. Methane is the major compound
of natural gas, and can be used as a fuel gas for SOFCs. For
the methane reformation process, water vapour was added to
the fuel gas. The study was performed with three different
gas mixtures consisting of methane, hydrogen, water and
argon (in ml/min.): 330–0–670–0, 280–125–560–60, and
235–235–470–60. These mixtures correspond to the
situation of no external reformation of methane and two
different pre-reformation situations. Experiments were
performed between 650 and 800 -C.
Table 5
Current density (A/cm2, at 700 mV) and area-specific resistance (mV cm2) of L58S
Temperature (-C) Current density (A/cm2, 700 mV) H2–H2O flow (ml
1000–30 1000–333 1000–1
800 1.27T0.10 1.24T0.17 1.23T0.
750 1.11T0.09 0.92T0.03 0.83T0.700 0.73T0.06 0.66T0.01 0.60T0.
650 0.47T0.04 0.41T0.04 0.35T0.
* Powder from ECN.
Fig. 9 shows as an example the current–voltage curves
as a function of temperature with 33 vol.% CH4 and 67
vol.% H2O. With decreasing temperature, the current
density at 700 mV was significantly reduced; from a
calculated average value of 0.88T0.05 A/cm2 at 800 -C to
0.29T0.03 A/cm2 at 650 -C (see also Table 6). A similar
behaviour was found for the other two methane-containing
gas compositions. These lower values of the current
densities and open circuit voltages (OCVs) can be explained
by the high amount of water resulting in a higher oxygen
partial pressure.
Table 6 also shows the average current densities and
area-specific resistances as a function of the gas composi-
tion. It can be concluded that with increasing H2 concen-
tration, and thus decreasing CH4 and H2O concentration, the
electrochemical performance was only slightly improved.
For example, the average current density at 700 mV and
calculated for 750 -C increased from 0.76T0.10 A/cm2 with
a 330–0–670–0 CH4–H2–H2O–Ar gas mixture to
0.86T0.07 A/cm2 with a 235–235–470–60 gas mixture,
whereas the area-specific resistance decreased from 412T59to 314T17 mV cm2.
It is worth noting that the oxygen partial pressure of a
fuel gas including methane with 67 vol.% H2O corresponds
to that of a gas mixture of hydrogen with about 33 vol.%
H2O. Due to this relatively high oxygen partial pressure a
lower OCV was reached: for example, the OCV at 800 -Cfor the methane–water mixture was about 970 mV. In the
case of hydrogen–water mixtures, the OCV for H2-25
CF-based single cells* as a function of the temperature and gas composition
/min) Area-specific resistance (mV cm2) H2–H2O flow (ml/min)
000 1000–30 1000–333 1000–1000
20 239T2 236T3 237T4
03 283T3 302T4 288T401 444T7 432T7 405T6
02 680T12 700T18 721T14
Table 6
Current density (A/cm2, at 700 mV) and area-specific resistance (mV cm2) of L58SCF-based single cells* as a function of the temperature and gas composition
Temp. (-C) Current density (A/cm2, 700 mV)
CH4–H2–H2O–Ar flow (ml/min)
Area-specific resistance (mV cm2)
CH4–H2–H2O–Ar flow (ml/min)
330–0–670–0 280–125–560–60 235–235–470–60 330–0–670–0 280–125–560–60 235–235–470–60
800 0.88T0.05 1.01T0.08 1.05T0.10 283T5 287T14 278T3
750 0.76T0.10 0.81T0.09 0.86T0.07 412T59 337T29 314T17700 0.56T0.02 0.57T0.03 0.59T0.03 597T107 499T24 467T36
650 0.29T0.03 0.30T0.03 0.32T0.04 967T55 853T79 826T51
* Powder from ECN.
A. Mai et al. / Solid State Ionics 176 (2005) 1341–1350 1349
vol.% H2O and H2-50 vol.% H2O was 990 and 910 mV,
respectively. From these data it can be concluded that the
OCV of the methane–water mixture corresponds more or
less to that of hydrogen including a relatively high
concentration of water.
3.5.3. Long-term behaviour
Endurance tests regarding the long-term electrochemical
behaviour of two SOFCs with L58SCF-type cathodes (one
based on the powder from ECN (The Netherlands), the other
one on spray-dried powder) were made for a period of 5200
and 2000 h, respectively (Fig. 10). Both single cells were
tested at 750 -C under a constant load of 0.5 A/cm2 with
1000 ml/min H2 (3 vol.% H2O) as the fuel gas. For the first
cell an average degradation rate of about 1.5% per 1000 h
was measured in the first 2200 h, which decreased to 1.0%
for the last 2600 h. For the second cell the degradation was
about 0.9% per 1000 h. This is rather high for single cell
measurements. Possible reasons for the degradation include
further SrZrO3 formation and coarsening of the micro-
structure due to sintering effects. Further work is in progress
to determine the reason for the loss in performance.
4. Conclusions
The applicability of mixed ionic-electronic conductive
perovskites as SOFC cathodes was compared with state-of-
0 1000 2000 3000 4000 50000.50
0.75
1.00
1.25
Cel
l vol
tage
(V
)
Time (h)
power blackout
Cel
l vol
tage
(V
)
Fig. 10. Long-term behaviour of two L58SCF-based cells (at 750 -C and a load
spray-dried L58SCF.
the-art LSM/YSZ cathodes. Although there are problems
with regard to chemical stability and thermal expansion,
some of the materials proved to be superior in electro-
chemical performance.
For a meaningful comparison of the electrochemical
performance of the cathode materials, the microstructure of
the samples has to be similar. This can be achieved by
applying different sintering temperatures adjusted for each of
the materials. Care has to be taken, however, that the TEC is
not exceptionally high (>18�10�6 K�1). Furthermore, the
materials should not show a high tendency to Sr depletion.
The performance of some types of perovskites was
significantly higher than that of Ftraditional_ LSM/YSZ
cathodes. Especially the LSCFs with 40% Sr on the A-site
and an A-site deficiency (L55SCF, L58SCF) showed high
power densities, in particular at low temperatures (at 0.7 V:
1.23 W/cm2 at 800 -C, 1.0 W/cm2 at 750 -C and 0.7 W/cm2
at 700 -C). This is nearly twice the power density of LSM/
YSZ cells.
It was shown that for L58SCF cathodes long-term
stability is still an issue to be considered, while the reasons
have to be investigated further.
Acknowledgements
The authors thank the staff of the IWV department at
Forschungszentrum Julich for processing the anode sub-
0 1000 20000.50
0.75
1.00
1.25
Time (h)
of 0.5 A/cm2). Left: cell based on powder from ECN. Right: cell based on
A. Mai et al. / Solid State Ionics 176 (2005) 1341–13501350
strates and electrolyte layers, namely Mr. Blaß, Mr. F. J.
Dias, Ms. M.-T. Gerhards, Mr. M. Kampel, and Ms. G.
Klein, and for synthesis of some of the cathode materials in
particular Mr. W. Jungen, Mr. W. Herzhof and Ms. K.
Portulidou, for SEM investigations Dr. D. Sebold, for the
SIMS analyses Dr. U. Breuer and Ms. A. Scholl, and for
performing the electrochemical measurements, in particular
Ms. C. Tropartz, Ms. B. Rowekamp, and Mr. H. Wese-
meyer. In addition, thanks go to Dr. G. Rietveld from ECN,
Petten, The Netherlands, for providing La0.58Sr0.4Fe0.8Co0.2O3� d powder and to Prof. H. Ullmann (TU Dresden)
and Dr. N.E. Trofimenko (FhG-IKTS, Dresden) for provid-
ing the CSCF powder.
The work was carried out in the networking project
‘‘Renewable Energies’’ under contract no. 01SF0039 and
financial support from the German Federal Ministry of
Science and Education (BMBF) is gratefully acknowledged.
References
[1] T. Lipman, D. Sperling, W. Vielstich, A. Lamm, H.A. Gasteiger
(Eds.), Handbook of Fuel Cells, vol. 4, John Wiley and Sons Ltd.,
Chichester, UK, 2003, pp. 1318–1328.
[2] D. Stover, H.P. Buchkremer, J.P.P. Huijsmans, in: W. Vielstich, A.
Lamm, H.A. Gasteiger (Eds.), Handbook of Fuel Cells, Volume 4:
Fuel Cell Technology and Applications Part 2, John Wiley and Sons
Ltd., Chichester, UK, 2003, pp. 1013–1031.
[3] M. Lang, T. Franco, R. Henne, S. Schaper, G. Schiller, in: A.J.
McEvoy (Ed.), Proceedings of the 4th European Solid Oxide Fuel Cell
Forum, European Fuel Cell Forum, Oberrohrdorf, Switzerland, 2000,
p. 231.
[4] B.C.H. Steele, Solid State Ionics 134 (2000) 3–20.
[5] A. Ahmad-Khanlou, F. Tietz, I.C. Vinke, D. Stover, in: H. Yokokawa,
S.C. Singhal (Eds.), Procceedings of the 7th Int. Symposium on Solid
Oxide Fuel Cells (SOFC VII), The Electrochemical Society Proceed-
ings, Pennington, NJ, 2001, p. 476 (PV 2001-16).
[6] M. Sahibzada, B.C.H. Steele, K. Zheng, R.A. Rudkin, I.S. Metcalfe,
Catal. Today 38 (1997) 459–466.
[7] Y. Teraoka, H.M. Zhang, S. Furukawa, N. Yamazoe, Chem. Lett.
(1985) 1743–1746.
[8] B.C.H. Steele, S. Carter, J. Kajda, I. Kontoulis, J.A. Kilner, in: F.
Grosz, P. Zegers, S.C. Singhal, O. Yamamoto (Eds.), Proceedings of
the 2nd Int. Symposium on Solid Oxide Fuel Cells, Commission of the
European Communities, Luxembourg, 1991, pp. 517–525.
[9] V.V. Kharton, A.A. Yaremchenko, E.N. Naumovich, J. Solid State
Electrochem. 3 (1999) 303.
[10] H.Y. Tu, Y. Takeda, N. Imanishi, O. Yamamoto, Solid State Ionics 117
(1999) 277.
[11] G. Stochniol, A. Gupta, A. Naoumidis, D. Stover, in: U. Stimming,
S.C. Singhal, H. Tagawa, W. Lehnert (Eds.), Proceedings of the 5th
Int. Symposium on Solid Oxide Fuel Cells (SOFC-V), The Electro-
chemical Society Proceedings, Pennington, NJ, 1997, pp. 888–896
(PV 97-40).
[12] A. Petric, P. Huang, F. Tietz, Solid State Ionics 135 (2000) 537.
[13] H. Uchida, S. Arisaka, M. Watanabe, Electrochem. Solid-State Lett. 2
(1999) 428–430.
[14] A. Tsoga, A. Gupta, A. Naoumidis, P. Nikolopoulos, Acta Mater. 48
(2000) 4709.
[15] P. Kontouros, R. Forthmann, A. Naoumidis, G. Stochiniol, E.
Syskakis, Ionics 1 (1995) 40.
[16] M.P. Pechini, United States Patent Office, Patent no. 3,330,697
(1967).
[17] D. Simwonis, A. Naoumidis, F.J. Dias, J. Linke, A. Moropoulou,
J. Mater. Res. 12 (1997) 1508–1518.
[18] D. Stover, H.P. Buchkremer, F. Tietz, N.H. Menzler, in: J. Huijsmans
(Ed.), Proceedings of the 5th European Solid Oxide Fuel Cell Forum,
European Fuel Cell Forum, Oberrohrdorf, Switzerland, 2002, pp. 1–9.
[19] F. Tietz, Ionics 5 (1999) 129–139.
[20] L.W. Tai, M.M. Nasrallah, H.U. Anderson, D.M. Sparlin, S.R. Sehlin,
Solid State Ionics 76 (1995) 259–271.
[21] H. Ullmann, N. Trofimenko, F. Tietz, D. Stover, A. Ahmad-Khanlou,
Solid State Ionics 138 (2000) 79–90.
[22] T. Ishihara, T. Kudo, H. Matsuda, Y. Takita, J. Electrochem. Soc. 142
(1995) 1519–1524.
[23] S.P. Simner, J.F. Bonnett, N.L. Canfield, K.D. Meinhardt, J.P. Shelton,
V.L. Sprenkle, J.W. Stevenson, J. Power Sources 113 (2002) 1–10.
[24] L. Kindermann, D. Das, H. Nickel, K. Hilpert, Solid State Ionics 89
(1996) 215–220.
[25] H.U. Anderson, Solid State Ionics 52 (1992) 33–41.
[26] G. Stochniol, E. Syskakis, A. Naoumidis, J. Am. Ceram. Soc. 78
(1995) 929–932.
[27] F. Riza, Ch. Ftikos, F. Tietz, W. Fischer, J. Eur. Ceram. Soc. 21 (2001)
1769–1773.
[28] A. Mai, ‘‘Katalytische und elektochemische Eigenschaften von eisen-
und kobalthaltigen Perowskiten als Kathoden fur die oxidkeramische
Brennstoffzelle (SOFC)’’, PhD thesis, Ruhr-Universitat Bochum
(2004).
[29] J. Misuzaki, I. Yasuda, J. Shimayoma, S. Yamauchi, K. Fueki,
J. Electrochem. Soc. 140 (1993) 467–471.
[30] J.A. Kilner, R.A. De Souza, I.C. Fullarton, Solid State Ionics 86-88
(1996) 703–709.
[31] A. Ahmad-Khanlou, ‘‘Alternative Werkstoffe fur Komponenten der
Hochtemperatur-Brennstoffzelle (SOFC) zur Herabsetzung der
Betriebstemperatur’’, PhD thesis, Ruhr-Universitat Bochum (2000).
[32] L. Qiu, T. Ichikawa, A. Hirano, N. Imanishi, Y. Takeda, Solid State
Ionics 158 (2003) 55–65.
[33] N.E. Trofimenko, H. Ullmann, J. Eur. Chem. Soc. 20 (2000)
1241–1250.
[34] M.T. Colomer, B.C.H. Steele, J.A. Kilner, Solid State Ionics 147
(2002) 41–48.
