AG-perovskite cermets for thin film solid oxide fuel cell air-electrode applications

An air electrode for a solid oxide fuel cell having high conductivity and low interfacial resistance is disclosed. Said air electrode is composed of an Ag-perovskite cermet.

BRIEF DESCRIPTION OF THE INVENTION 
This invention relates generally to thin film solid oxide fuel cells 
(SOFC's) and more particularly to the Ag-perovskite cermet air electrodes 
in SOFC's. 
BACKGROUND OF THE INVENTION 
Solid oxide fuel cells generally comprise a solid electrolyte such as 
yttria stabilized zirconia (YSZ) with a positive air electrode on one 
surface of the electrolyte and a negative fuel electrode such as a cermet 
of nickel and stabilized zirconia on the other surface of the electrolyte. 
Results on solid oxide fuel cells (SOFC's) employing thin (.about.10 .mu.m) 
Y-stabilized zirconia (YSZ) electrolytes [1,2] have recently been 
reported. One of the main advantages of these thin-film-electrolyte cells 
is that electrolyte ohmic resistance is small even at reduced 
(600-800.degree. C.) operating temperatures T.sub.c [3]. Reduced SOFC 
temperatures should ease materials and processing problems associated with 
interconnection and gas sealing in SOFC stacks [4]. However, as T.sub.c is 
lowered below the typical value of .apprxeq.1000.degree. C., the rates of 
thermally-activated electrochemical reactions decrease, increasing 
electrode interfacial resistance r.sub.i and limiting the SOFC current 
density. In addition, the resistivity .rho. of perovskite air electrode 
materials increases with decreasing temperature. 
Two general approaches have been used to reduce r.sub.i. First, the 
electrode and/or electrolyte compositions can be changed to provide 
improved catalytic performance. For example, (La,Sr)CoO.sub.3 (LSC) air 
electrodes on YSZ have been shown to provide improved performance over 
(La,Sr)MnO.sub.3 (LSM). The addition of a thin interfacial layer of a 
second electrolyte, such as Y-doped Bi.sub.2 O.sub.3 (YSB), can also lower 
r.sub.i [5]. Second, standard electrode-electrolyte combinations can be 
processed so as to maximize he three-phase boundary length. 
Electrochemical vapor deposition has been used to deposit YSZ within 
porous LSM electrodes to increase the contact area. Another means for 
increasing the contact area is via the use of a cermet electrode, e.g. 
Pt-YSZ or Ag-YSZ, where the additional contact area is within the cermet. 
Cermet electrodes can also provide lower resistivity than ceramic air 
electrodes, especially at low temperatures. For Ag-YSZ cermets[6], for 
example, .rho. values were .apprxeq.30 times lower than those of LSC and 
LSM at 750.degree. C. There has been relatively little work done on 
cermets that contain noble metals and perovskite oxides, with only one 
preliminary report on Pt-LSM [7] for typical T.sub.c values of 
.apprxeq.1000.degree. C. We have observed that for SOFC's with 
Tc&lt;800.degree. C. Ag is an excellent candidate for cermets using 
perovskite oxide because of its good catalytic activity, high electrical 
conductivity, and relatively low cost. 
OBJECTS AND SUMMARY OF THE INVENTION 
It is an object of this invention to provide an inexpensive air electrode 
for SOFC's which has good catalytic activity and high electrical 
conductivity. 
It is another object of the invention to provide thin film Ag-perovskite 
cermets for use as positive air electrodes in SOFC's. 
Pursuant to the invention there is provided a sputtered Ag-perovskite thin 
film air electrode for SOFC's.

DESCRIPTION OF THE PREFERRED EMBODIMENT 
Cermet electrode films in accordance with this invention were deposited by 
dc magnetron sputtering in a cylindrical, 35 cm diameter, glass bell jar 
vacuum chamber that was described in more detail elsewhere [8]. The 
chamber was turbomolecular pumped and reached a base pressure of 
&lt;3.times.10.sup.-6 Torr when the liquid-nitrogen trap was filled. Two 
spaced planar magnetron sputter sources 11, 12 FIG. 1 were used, one with 
a 99.95% pure, 5-cm-diameter Ag target and the other with the LSC (or LSM) 
target. The substrate plane 13 was parallel to the target surface plane, 
at a distance of 7.6 cm. The distance between the target centers was 10 
cm. Substrates were placed along a line above the centers of the 
sputtering targets, allowing a wide range of cermet compositions to be 
deposited in a single deposition, while the cermet compositions were 
constant to within 5% over the 1 cm sample size. 
(La.sub.0.7 Sr.sub.0.3)CoO.sub.3 and (La.sub.0.7 Sr.sub.0.3)MnO.sub.3 
perovskite targets were synthesized by standard ceramic processing or 
purchased from SSC, Inc. In the former case, starting powder materials 
were La.sub.2 O.sub.3 (99.99%), SrCO.sub.3 (99.9%) and Co.sub.3 O.sub.4 
(99.5%). The powders were weighed in the desired ratio, dissolved in 
acetone solution and mixed for about 1 h. Finally, the mixed powders were 
cold-pressed at 100 MPa and sintered at 1150.degree. C. in air for 30h 
into discs 5 cm diameter and .apprxeq.3 mm in thickness. The formation of 
the perovskite structure was confirmed by means of x-ray diffraction (XRD) 
analysis. Energy dispersive x-ray analysis (EDX) showed the LSC and LSM 
targets to have the desired metal compositions, i.e. 30 mol % Sr on the A 
site and stoichiometry between the A and B sites. The commercial 
perovskite powders were cold pressed at 25 MPa and sintered in air at 
1150.degree. C. for 10 h. No difference was observed in the results for 
the two different methods of target preparation. 
Films were deposited onto silica glass substrates except in cases where 
high temperature (.gtoreq.600.degree. C.) experiments were to be carried 
out, in which case dense alumina or YSB (25 mol % Y.sub.2 O.sub.3 
substrates were used. Alumina and YSB substrates were polished using 1 
.mu.m diamond-embedded paste. The polished substrates were ultrasonically 
cleaned in 1 H.sub.3 PO.sub.4 :1 H.sub.2 O solution to remove surface 
contamination. Prior to film deposition, both kinds of substrates were 
successively rinsed in de-ionized water, ultrasonically cleaned in acetone 
and finally dried with dry nitrogen. The substrates were not intentionally 
heated, but the substrate temperature rose to .apprxeq.70.degree. C. 
during deposition under typical conditions. 
The sputtering gas was 99.999% pure argon. The argon pressure P was 10 
mTorr, corresponding to an argon flow rate of 114 ml/min. For the Ag 
target, the applied voltage was 300 V and the current was 0.05 A, while 
the LSC (or LSM) target was operated at 200 V and 0.5 A, respectively. 
Prior to film deposition, the targets were sputter cleaned for 10 mins 
with a shutter covering the substrates. New perovskite targets were also 
sputtered for several hours to reach a steady state condition prior to 
depositing films. While cracks formed in the LSC and LSM targets after the 
initial sputtering period, the discharge conditions, sputtering rates, and 
film purity remained unaffected. Table 1 summarizes the sputtering 
conditions. Some of the sputter-deposited films were annealed in air at 
500-600.degree. C. for 10h prior to characterization. 
TABLE 1 
______________________________________ 
Ag-LSC-Co-Sputtering Conditions 
Targets Ag LSC 
______________________________________ 
Voltage (V) 300 200 
Current (A) 0.05 0.5 
Sputter Gas Ar(10 mTorr) 
Substrate Temperature 70.degree. C. 
Deposition Rate 0.7-1.6 .mu.m/h 
______________________________________ 
The film thicknesses were 0.5-5 .mu.m, as measured by scanning a 
profilometer stylus across a step edge created by masking the substrate 
during deposition. Film chemical compositions were measured using energy 
dispersive x-ray analysis (EDX). Film stresses were determined by 
measuring the curvature of the substrate before and after film deposition 
with a microstylus profilometer with a scan range of 1 cm. Morphology was 
analyzed using scanning electron microscopy (SEM) with chemical mapping. 
Crystal structure was determined using x-ray diffraction (XRD) in the 
.theta.-2.theta. geometry with Cu K.alpha. radiation (40 kV-20 mA). 
Electrical resistivities were measured using a van der Pauw geometry with 
gold or silver wires connected to the sample using silver paste. Electrode 
polarization curves were measured using a potentiostat (EG&G model 273) 
connected to a personal computer for data storage and analysis. 
Interfacial resistance of Ag-perovskite electrodes on (Y.sub.2 
O.sub.3).sub.0.25 (Bi.sub.2 O.sub.3).sub.0.75 were studied using a 
Solartron SI 1260 Frequency Response Analyzer (FRA) over a frequency range 
100 kHz to 10 mHz using a .+-.10 mV excitation signal. Two sample 
geometries were used. First, the 0.5-1.0 .mu.m thick electrodes were 
deposited on both sides of a 2-mm-thick YSB ceramic. The area-specific 
resistance r.sub.i =R.sub.c A, where A is the electrode area and R.sub.e 
is the measured resistance. Second, measurements were carried out between 
two electrodes deposited on 3-.mu.m-thick YSB (25 mol % Y.sub.2 O.sub.3) 
electrolyte thin films on dense alumina substrates. The YSB was deposited 
by reactive co-sputtering from Y and Bi targets in Ar/O.sub.2 mixtures at 
a total pressure of 10 mTorr. The Ag-perovskite cermet electrodes (0.5-1.0 
.mu.m thick) were deposited through a mask that defined two symmetric 
rectangular regions (FIG. 2) on the YSB film, of width w=0.6-1.0 cm and 
length d.sub.2 =2.5 mm, separated by a distance d.sub.1 =0.4-0.6 mm. 
Comparisons of the two types of measurements gave excellent agreement. 
The film deposition rate from two targets agreed well with the 
superposition of the rates from individual targets, as shown previously 
for Ag-YSZ cermets [6]. The silver volume fraction f.sub.Ag as a function 
of substrate position x (FIG. 1) is shown in FIG. 3. The composition 
estimated from the superposition of the deposition rates of Ag, R.sub.Ag, 
and LSC, R.sub.LSC, using the expression f.sub.Ag (x)=R.sub.Ag 
(x)(R.sub.Ag (x)+R.sub.LSC (x)], is also plotted in FIG. 3. The curves 
agreed within experimental accuracy, indicating that there was little 
cross-contamination between the targets. EDX measurements showed that the 
metal composition of the deposited perovskite phase was stoichiometric 
within the experimental error (&lt;5%), and was not changed after annealing. 
Film densities were estimated by dividing the film mass, given by the mass 
difference before and after deposition, by the film volume, obtained from 
the film thickness and area. Because of the errors in these two 
measurements, the overall accuracy is 10%. For both pure Ag and cermet 
films with f.sub.Ag =0.3, the density .eta. relative to the theoretical 
value .eta..sub.th is .eta./.eta..sub.th .apprxeq.0.7. 
For stress measurements, the films were deposited on 200-.mu.m-thick glass 
substrates. The film stress S was related to the deflection .delta.f the 
substrate, measured between the substrate center and a distance D from the 
center, by the equation 
##EQU1## 
where E it the Young's modulus of the substrate (69 GPa), t.sub.s is the 
substrate thickness t.sub.f is the film thickness and v is Poisson's ratio 
for the substrate (0.16). S is plotted as a function of f.sub.Ag for 
Ag-LSC in FIG. 4. The films were in a compressive stress state. S was 
largest for put LSC, 3.5.times.10.sup.8 Pa, and decreased rapidly with 
increasing f.sub.Ag to 0.5.times.10.sup.8 Pa for pure Ag. The magnitude 
and sign of the stress are very similar to previous results for Ag-YSZ 
cermet films prepared using the same technique [6]. The lower stress in 
Ag-rich films was likely due to plastic deformation of Ag limiting the 
film stress to relatively low levels. 
FIGS. 5A-D show representative x-ray diffraction (XRD) spectra from both 
as-deposited and annealed 1-.mu.m-thick films with different f.sub.Ag. The 
as-deposited f.sub.Ag =0.5 cermet film (a) was partially amorphous as 
indicated by the lack of perovskite peaks, the weak broadened Ag peak 
intensities, and relatively large background intensity. The crystallinity 
of this sample improved after annealing at 500-600.degree. C. for 10h (b) 
as indicated by the sharper, more-intense Ag peaks and the appearance of 
LSC peaks. XRD results from f.sub.Ag =0.3(c) and f.sub.Ag =0(d) films 
after annealing are also shown. The relative Ag and perovskite peak 
intensities correlated well with the film composition. 
The Ag and perovskite distribution in the cermet films was assessed using 
SEM chemical imaging. FIG. 6A shows the chemical images of the 
spatially-resolved La, Sr, Co, and Ag EDX signals in an as-deposited 
f.sub.Ag =0.3 Ag-LSC film deposited on an alumina substrate. La, Sr, Co, 
and Ag were uniformly dispersed, on the scale observed by SEM 
(.apprxeq.0.1 .mu.m), in the as-deposited film. FIG. 6B shows the EDX 
chemical image of the same film after annealing at 750.degree. C. for 2h 
in air. The only change from the as-deposited film in FIG. 6A was the 
appearance of isolated Ag clusters. Segregation of the metal out of the 
cermet structure has also been observed for Ag-YSZ [6] and Ni-YSZ cermets 
after annealing. The segregation can be suppressed by the use of a pure 
LSC cap layer, as discussed below. 
The resistivity .rho. of 1-.mu.m-thick Ag-LSC films on alumina substrates 
versus f.sub.Ag, measured at temperature T=750.degree. C. in air, is shown 
in FIG. 7. The films had been annealed at 600.degree. C. in air for 10h 
prior to measurement, and .rho. values were stable during the measurement. 
.rho. decreased with f.sub.Ag from 1.6.times.10.sup.-2 .OMEGA..cm for a 
pure perovskite film to 1.times.10.sub.-5 .OMEGA..cm for a pure Ag film. 
The .rho. variation with metal volume fraction was much stronger on the 
perovskite-rich side than that observed for Pt-LSM cermets [7]. However, 
the present resistivity values are in excellent agreement with the general 
effective media (GEM) theoretical predictions for cermets where both 
phases are conducting. The calculated cermet electrical conductivity 
.sigma. is given by 
##EQU2## 
where .sigma..sub.Ag and .sigma..sub.LSC are the conductivities of the two 
phases, A=(1-f.sub.c)/f.sub.c, f.sub.c is the critical Ag volume fraction 
for percolation conductivity, and p is a parameter usually in the range 
from 1.65-2.0. The solid line shown in FIG. 6 was obtained using f.sub.c 
=0.4 and p=2.0, which are both reasonable values. The calculated curve is 
in good agreement with the data. 
.sigma.T for annealed films with different Ag-LSC compositions is plotted 
versus 1000/T in FIG. 8. Straight lines are observed over the whole 
temperature range, and the slope is consistent with a small-polaron 
conduction model for the LSC-rich compositions. With increasing Ag 
content, the slope decreases and eventually reverses, exhibiting the 
decrease in .sigma. expected for metals with increasing T. 
Impedance spectroscopy studies of reactions at Ag-LSC electrodes on YSB 
were carried out as a function of T, oxygen partial pressure P, and 
f.sub.Ag. FIGS. 9A-D show complex impedance plots for different electrode 
compositions on YSB electrolytes, measured at T=750.degree. C. in air. The 
bulk arcs were omitted. The f.sub.Ag =0, 0.3, and 0.5 results showed a 
superposition of two arcs, indicating that there was more than one 
important rate-limiting step. The lower-frequency arc became more 
important with increasing f.sub.Ag. 
The electrode interfacial resistance r.sub.1, determined from the real-axis 
intercepts of the arcs as shown in FIGS. 9A-D, are plotted versus f.sub.Ag 
for both Ag-LSC ad Ag-LSM in FIG. 10. For Ag-LSC r.sub.i first decreased 
with increasing f.sub.Ag, reached a minimum near f.sub.Ag =0.3, and then 
increased rapidly to the value characteristic of pure Ag. The shape of the 
impedance arcs for Ag-LSM was similar to those shown for Ag-LSC. The 
r.sub.i value for pure LSM was almost a factor of two higher than that for 
pure LSC. Adding Ag to LSM decreased r.sub.i substantially, as shown for 
f.sub.Ag =0.3 in FIG. 10. 
FIG. 11 gives r.sub.i as a function of 1/T for Ag, LSC, and cermet 
electrodes with f.sub.Ag =0.3 and 0.5. The f.sub.Ag =0.3 composition 
showed the lowest interfacial resistance for T&gt;550.degree. C., while pure 
Ag had the lowest r.sub.i for T&lt;550.degree. C. The slope of logr.sub.i 
versus 1/T was largest for f.sub.Ag =0.3 (the corresponding activation 
energy was 105 kJ/mol), and was sightly smaller for both pure Ag, 
corresponding to 58 kJ/mol. The similar slopes for LSC and f.sub.Ag =0.3 
and 0.5 electrodes suggests that the rate-limiting steps were the same, 
although the f.sub.Ag =0.3 cermet exhibited a lower absolute r.sub.i rate. 
FIGS. 12A-D show the low-frequency impedance diagrams as a function of P 
for the f.sub.Ag =0.3 cermet. Two overlapping electrode arcs were observed 
in each case, with the lower-frequency arc dominating at low P and the 
higher-frequency arc dominating at high P. FIG. 13 shows r.sub.i versus P 
for Ag, LSC, and f.sub.Ag =0.3 electrodes obtained from data such as shown 
in FIGS. 12A-D. The Ag electrode exhibited a r.sub.i .alpha. P.sup.-0.5 
dependence, indicative of mass transport as the rate-limiting step. Both 
the LSC and cermet electrodes exhibited a P.sup.-0.5 power law at 
intermediate P, but showed indications of a larger exponent at lower P and 
a smaller exponent at higher P. For electrode operation in air, the 
primary rate-limiting process was thus mass transport, although there were 
indications that charge transfer played a role. The P-dependencies for the 
LSC and f.sub.Ag =0.3 electrodes were essentially the same, indicating 
similar reaction mechanisms in agreement with the temperature-dependent 
studies above. 
The overpotential .eta. of the f.sub.Ag =0.3 electrode on YSB was measured 
with a potentiostat at T=750.degree. C. in air. The low current resistance 
was very close to the value r.sub.i =0.3 .OMEGA.cm.sup.2 measured by 
impedance spectroscopy, but increased slightly with increasing current 
density to yield .eta..apprxeq.35 mV at 100 mA/cm.sup.2. 
To summarize the above results for T=750.degree. C. in air, the electrode 
resistances were limited primarily by mass transport, although charge 
transfer apparently became more important for the LSC-rich compositions. 
Two possible explanations can be proposed for the decrease in r.sub.i as 
Ag is added to LSC, as shown in FIG. 10. First, the combination of two 
materials with different rate-limiting processes may allow a 
"short-circuiting" of the limiting steps of each. For example, Ag may 
provide sites for efficient charge transfer, while LSC may provide rapid 
mass transport. Second, the effective three-phase boundary length in the 
cermet may be larger due either to a composition-dependent change in 
electrode morphology or the additional Ag-LSC interfaces in the cermet. 
Based on the above results, Ag-perovskite cermets with f.sub.Ag in the 
range of 0.2-0.6 provide high conductivity and low interfacial resistance. 
Cermets with volume fraction of silver of about 0.3 are preferable because 
they combine high conductivity with low interface resistance. 
This composition was thus tested in initial 200h stability tests in air at 
750.degree. C. Based upon prior experience with Ag-YSZ electrodes [5], the 
tests were conducted on electrodes that were coated with a 1 .mu.m thick, 
pure porous LSC layer in order to suppress Ag evaporation and segregation. 
The cap layer was effective in maintaining stable electrode morphology and 
composition, as indicated by SEM and EDX observations after annealing 
r.sub.i for 200 h at 750.degree. C. were examined by ERD and showed no 
evidence of the formation of interfacial phases. 
The poor thermal expansion match between these materials and the 
electrolyte is a potential problem, especially for Ag-LSC. There was no 
evidence of cracking or delamination of the Ag-perovskite films despite 
numerous thermal cycles between room temperature and 750.degree. C. This 
indicates that at a film thickness of 1 .mu.m, the Ag and LSC thermal 
expansion difference with alumina and YSB substrates was not a major 
problem. Ag-perovskite films with thickness in the range of 1.5 to 100 
.mu.m are useful; however, Ag-perovskite films with thickness in the range 
of 10.5 to 1.5 .mu.m are preferred, because that range minimizes the 
cracking problem. 
The resistivity of Ag-(La.sub.0.7 Sr.sub.0.3)CoO.sub.3 cermets decreased as 
the Ag volume fraction f.sub.Ag increased from 
.apprxeq.1.6.times.10.sup.-2 .OMEGA..cm for pure LSC to .apprxeq.10.sup.-5 
.OMEGA..cm for pure Ag at 740.degree. C. The f.sub.Ag =0.3 cermet had a 
resistivity factor of =20 times less than pure LSC. The cermet electrodes 
should thus allow either a decrease in electrode resistance, important to 
reduce ohmic polarization in cell geometries with long current paths, or a 
decrease in electrode thickness to reduce materials costs. The interfacial 
resistance r.sub.i of the cermet electrodes with f.sub.Ag =0.3 on YSB 
electrolytes was 0.3 .OMEGA..cm.sup.2 at 750.degree. C. in air, lower than 
either pure LSC or pure Ag, and was stable during initial long-term 
testing. Ag-LSM electrodes also exhibited lower r.sub.i than either pure 
Ag or pure LSM on YSB electrolytes. These cermet materials thus have 
desirable properties for medium-temperature SOFC applications; including 
low resistivity, excellent catalytic performance, and good stability. 
Although the cermet material has been described particularly with respect 
to SOFC's, the cermet is also useful for other solid-state electrochemical 
devices, e.g., sensors and electrolyzers. 
REFERENCES 
[1] L. S. Wang and S. A. Barnett, Solid State Ionics 61 (1993) 273. 
[2] N. Q. Minh, T. R. Armstrong, J. R. Esopa, J. V. Guiheen, C. R. Home, 
and J. J. VanAckeren, in: Third Int. Symp. Solid Oxide Fuel Cells, eds. S. 
C. Singhal and H. Iwahara Electrochem. Soc., Pennington, 1993) p. 801. 
[3] S. A. Barnett, Energy 15 (1990) 1. 
[4] K. Krist and J. D. Wright, in: Proc. Third Int. Symp. Solid Oxide Fuel 
Cells, eds. S. C. Singhal and H. Iwahara (Electrochem. Soc., Pennington, 
1993) p. 782. 
[5] L. S. Wang and S. A. Barnett, J. Electrochem. Soc. 139 (1992) 89. 
[6] L. S. Wang and S. A. Barnett, J. Electrochem. Soc. 139 (1992) 1134. 
[7] A. Tsunoda, T. Yoshida, and S. Sakurada, in: Proc. First Int. Symp. 
Solid Oxide Fuel Cells, ed. S. C. Singhal (Electrochem. Soc., Pennington, 
1989) p. 204. 
[8] E. S. Thiele, L. S. Wang, T. O. Mason, and S. A. Barnett, J. Vac. Sci. 
Technol. A9 (1991) 3054.