Journal of The Electrochemical Society, 161 (6) F717-F723 (2014) F7170013-4651/2014/161(6)/F717/7/$31.00 © The Electrochemical Society

Infiltrated Lanthanum Nickelate Cathodes for Use withBaCe0.2Zr0.7Y0.1O3−δ Proton Conducting ElectrolytesSean M. Babiniec,z Sandrine Ricote, and Neal P. Sullivan∗

Department of Mechanical Engineering, Colorado School of Mines, Golden, Colorado 80401, USA

Symmetric cells (electrode/electrolyte/electrode) were prepared using BaCe0.2Zr0.7Y0.1O3−δ (BCZY27) as a proton conductingelectrolyte. Porous BCZY27 backbones were sintered on each side of the electrolyte. Lanthanum nickelate, La2NiO4+δ (LN), wasinfiltrated into the porous BCZY27 backbone to act as the electrode. Pure phase LN was formed after annealing in air at 850◦C for4 h. Electron microscopy was used to observe LN nano-particles (∼50 nm) infiltrated throughout the backbone. Impedance spectrawere recorded at 500, 600, and 700◦C with four oxygen partial pressures and four steam partial pressures. Four arcs were observedin the spectra corresponding to: BCZY27 backbone grain-boundary resistance, charge-transfer, oxygen dissociation/adsorption, andgas transport. Total area specific resistances (ASRs) were found to decrease with increasing oxygen partial pressure. Total cathodeASRs of 161.8, 29.3, and 3.7 �·cm2 were measured in 3% humidified air at 500, 600, and 700◦C respectively.© 2014 The Electrochemical Society. [DOI: 10.1149/2.037406jes] All rights reserved.

Manuscript submitted February 26, 2014; revised manuscript received March 28, 2014. Published April 24, 2014.

High-temperature proton conductors (HTPCs) have been studiedin detail as electrolyte materials for fuel cells for over two decades.1-5

Protonic-ceramic fuel cells (PCFCs) are of interest due to the poten-tial of lower operating temperature (∼600◦C) as compared to fuelcells utilizing traditional oxygen-ion conducting electrolytes.6 Addi-tionally, the electrochemical production of steam in PCFCs occurs onthe cathode side, eliminating fuel dilution on the anode side. Manyadvances have been made in the development of proton conductingmembranes, especially in barium cerates and barium zirconates.4,5,7

However, interest in work on cathode materials for such membranesis relatively new.

It is imperative to increase the three-phase boundary (TPB) den-sity of the cathode to lower electrochemical resistance in the cathodeand increase cell performance. In traditional oxygen-ion solid-oxidefuel cells (SOFCs), this is done by utilizing materials with mixedoxygen-ion/electron conductivity (MIECs). However, the presence ofproton conductivity in PCFCs as opposed to oxygen-ion conductivitychanges the cathode requirements for three-phase boundary (TPB)length. MIECs (oxygen-ion/electron) are not suited for PCFC use,as proton conductivity is required to extend the TPB length. He et al.proposed a reaction model for PCFC cathodes and applied it to EIS re-sults of Sm0.5Sr0.5CoO3−δ (SSC)-BaCe0.8Sm0.2O3−δ (BCS) cathodes.8

In their study, they maintain that the half-reaction at the cathode of aPCFC is:

4OH•O + O2(g) + 4e′ ⇀↽ 2H2O(g) + 4O×

O [1]

The cathode half-reaction for an SOFC is:

O2(g) + 4e′ + V••O

⇀↽ 2O×O [2]

It is evident that the difference in cathode reactions between PCFCsand SOFCs results in different cathode processes and material require-ments. Each process is said to have an oxygen and steam dependency,as stated in Equation 3:8

Rpi ∝ (pO2)−mi (pH2O)−ni [3]

where Rp is the polarization resistance, and i refers to a specificprocess.

These key differences between SOFC and PCFC cathodes requiredevelopment of TPB densities to be pursued both architecturally andthrough materials advancement. Figure 1 illustrates the different op-tions in architecture in PCFC cathodes. If no protonic conductiv-ity is present in the cathode, the TPB is limited to the small areawhere the electrode material meets the electrolyte and the gas phase(Figure 1(a)). To further increase TPB densities, ceramic/ceramiccomposites (cercers) can be utilized that combine both an electronicphase and a protonic phase (Figure 1(b)). Figure 1(c) illustrates an

∗Electrochemical Society Active Member.zE-mail: sbabinie@mines.edu

electronically conductive nano-catalyst infiltrated into a porous pro-ton conducting backbone. The small particle size of the nano-catalystincreases the area of contact between the protonic phase and the elec-tronic phase, increasing the TPB density. However, the nano-particlesmust percolate to provide electrons from the current collection to theelectrochemical sites.

Previous studies have used two-phase cercer composites (pro-tonic ceramic phase/MIEC phase) to increase the TPB length, buthigh sintering temperatures caused cation interdiffusion between thetwo phases.9,10 An infiltrated-cathode technique can be used to mit-igate cation diffusion, and has been used for SOFCs for quite sometime.11-14 The small particle size (<100 nm) of the infiltrated cat-alyst increases the surface area and the three phase density whencompared to cercer cathodes (Figure 1). Ricote et al. demonstratedthis approach by fabricating symmetric cells consisting of infiltratedLaCoO3 nano-particles into a BaCe0.2Zr0.7Y0.1O3−δ (BCZY27) porousbackbone with a BCZY27 electrolyte.15 They measured a polarizationresistance of 0.14 �·cm2 at 600◦C in 3% humidified air, one of thelowest observed for a PCFC cathode.

In 2010, Dailly et al. illustrated the value of A2MO4-type ox-ides as a potential cathode material for PCFCs.16,17 They noted thattraditional perovskite structures had little to no water uptake underthermogravimetric analysis (TGA) measurements, while materials ofthe Ruddlesden-Popper family (A2MO4) showed some water uptake,most notably La2NiO4+δ (LN). This uptake of water suggests high-temperature proton incorporation into the material, implying protonic-conductivity and/or steam decomposition activity, which are both de-sirable for fuel cell and electrolysis cell applications. This encourag-ing result motivates further study of A2MO4-type materials for PCFCcathode applications.

In this work, the performance of infiltrated La2NiO4+δ was studiedin a symmetric-cell configuration. Dense BaCe0.2Zr0.7Y0.1O3−δ sub-strates (1 mm thick) were fabricated with ∼40 μm porous BCZY27backbones on each side. La2NiO4+δ was infiltrated into the backboneuntil percolating at a loading of 7.3 wt%. An identical “control” cellwas also fabricated without nano-catalyst infiltrated into the back-bone. Electrochemical impedance spectroscopy (EIS) experimentswere conducted in a variety of environments to investigate the po-larization characteristics of the cells. X-ray diffraction (XRD), scan-ning electron microscopy (SEM), and energy dispersive spectroscopy(EDS) were also used to characterize the cells before and after testing.

Experimental

Preparation of symmetric cells.— Dense BCZY27 substrates weremade using solid-state reactive sintering, previously described in theliterature.18-21 During solid-state reactive sintering, phase formationfrom precursor oxides and particle sintering occur simultaneously.To assist in this reaction, 1 wt% NiO was added to the precursor

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F718 Journal of The Electrochemical Society, 161 (6) F717-F723 (2014)

H2O(g)O2(g)

H+

e-O2(g)

H2O(g)

BCZY27H+

O2(g)

H2O(g)

BCZY27

H+ H+

e-

H+BCZY27

H+

e-

H+

a) b) c)

Figure 1. Illustration comparing three-phase boundary geometry of a) tradi-tional MIEC on the BCZY27 proton-conductor, b) a MIEC/BCZY27 cercercomposite, c) infiltrated BCZY27 with nano-catalysts. TPBs are shown in red.

powders as a sintering aid. Although the exact sintering mechanism isnot fully understood, many studies offer possible mechanisms.20,22,23

This phenomenon reduces the number of processing steps requiredin making a ceramic part. Precursor powders of barium carbonate(99.95% metals basis), cerium oxide (99.5% REO), zirconium oxide(99.7% metals basis), and yttrium oxide (99.9% REO), all obtainedfrom Alpha-Aesar, were mixed in stoichiometric ratios and ball milledin water with 1 wt% NiO for 24 hours. 3 wt% polyethylene glycol wasadded as a binder, and the mixture was dried and sieved to -100 meshto remove large agglomerates. Pellets (dia = 20 mm) were made byuni-axially pressing the powder to 35 MPa and sintering in air to1535◦C with a 10 hour dwell. The resulting substrate was hand-machined on a polishing wheel to a thickness of 1 mm.

In fabrication of a porous backbone, small particle sizes and highporosity are desired. Solid-state reactive sintering typically resultsin large grains (∼5–30 μm).21 Such grain sizes are undesirable in acathode as catalytic surface area is reduced. Therefore, precalcinedBCZY27 was prepared by calcining the precursor powders at 1350◦Cfor 24 hours in air. NiO was added to the BCZY27 powders at 4 wt%as a sintering aid to promote sintering at lower temperatures. A pastevehicle was developed by mixing 3 wt% ethyl-cellulose to α-terpineol.Powders were mixed into the vehicle at a 5:3 oxide:vehicle ratio. Toimprove adhesion without increasing the sintering temperature andcompromising porosity, a non-percolating, low solids loading NiOsolution was dripped onto both sides of the substrate and allowedto dry. This sintering aid layer is crucial to the success of backboneadhesion to the electrolyte. The BCZY27 backbone paste was screenprinted to both sides of the substrate over the NiO adhesion layer. Thebackbones were sintered at 1375◦C for 5 hours in air.

To create the infiltration solution, lanthanum(III) nitrate (Alpha-Aesar, 99.9985% metals basis) and nickel(II) nitrate (Alpha-Aesar,99.9% REO) were dissolved in water in Stoichiometric ratios to aconcentration of 1 M. After the nitrates were dissolved, citric acidwas introduced at a 1:1 molar ratio. The cation solution was drippedinto the backbone 10 times, with a calcine at 450◦C between eachinfiltration. After the 10 infiltrations, a final calcine was performed at850◦C for 4 hours in air. One cell was fabricated without the lanthanumnickelate catalyst infiltration. The cell with the infiltration is termedthe “LN” cell, and the one without any catalyst is termed the “control”cell.

Characterization of the symmetric cells.— Crystallographic iden-tification was accomplished using a PANalytical X’Pert X-ray diffrac-tometer utilizing Cu K-α radiation.

Cross-sectional micrographs of fractured and polished sampleswere collected using a JEOL JSM-7000 field emission scanning elec-tron microscope (FESEM). High-resolution imaging was used to de-termine particle sizes, layer thickness, porosity, and layer adhesion,as well as identify infiltrated phases. Energy dispersive spectroscopy(EDS) was used to create element maps to investigate migration ofcations. A thin layer of gold was sputtered onto each sample to miti-gate charging effects.

For AC impedance measurements, porous Ag current collectors(ESL) were painted onto the electrodes of both samples and sin-tered in-situ at 700◦C before testing. Silver wires were connected in

Table I. Symmetric cell EIS experimental matrix.

pO2 (atm) 0.01 0.21 0.50 1.00

pH2O (atm)0.00 X0.03 X X X X0.10 X0.20 X

a four-point configuration using silver paste. Impedance spectra wererecorded from 500 to 700◦C in increments of 100◦C in four differentoxygen partial pressures and four different steam partial pressures.Oxygen partial pressure was controlled by mixing oxygen and argon.Addition of steam was accomplished by flowing gas through tem-perature controlled bubblers. All measurements were performed atequilibrium and under zero-DC current conditions. Table I shows thetesting matrix for the experiments. The spectra were recorded with aGamry Reference 600 potentiometer in the frequency range of 1 MHzto 5 mHz, with an amplitude of 25 mV and no DC offset.

To ensure the statistical validity of the measured data, a Kramers-Kronig (K-K) test was performed.24 The EIS data was fit using theZsimpwin software (EChem Software: echemsw.com).

Results and Discussion

Phase and microstructural characterization.— Figure 2 shows theXRD pattern of the BCZY27 backbone infiltrated with LN. Compar-ing the data to the BCZY27 and LN ICDD cards shows that there is nounexpected or missing peak. The NiO 1 and NiO 2 lines correspondto the 91 and 100 peak intensities of NiO, and La2O3 1 correspondsto the 100 peak intensity of La2O3. The absence of peaks in these lo-cations suggests that there is no significant catalyst precursor materialremaining. The slight peak shifting observed is due to the topographyof the backbone.

To test the stability of LN in high-steam environments, La2NiO4+δ

powders were calcined from the nitrate solution at 1100◦C andplaced in a 15% steam environment in air at 700◦C for five days.Figure 3 shows the XRD pattern of the LN after exposure to highsteam environments. The pattern matches the ICDD data points, andthere is no evidence of additional phases. This implies that the LNmaterial is stable under such high-temperature, high-steam conditions.

Electron micrographs of LN-infiltrated cathodes are shown inFigure 4, both before and after testing. Pre-test micrographs (Figs.4(a) and 4(b)) reveal ∼ 50 nm nano-particles on the surface of thebackbone. Backbone particles are ∼ 0.5 - 1 μm in diameter. Toconfirm that these particles are indeed lanthanum nickelate and notBCZY27 surface texture, EDS was performed on the sample anda lanthanum element map was made. A section was chosen at theelectrolyte-electrode interface so that it will be clear if any significant

Figure 2. XRD pattern of La2NiO4+δ infiltrated into BCZY27 backbone.Square symbol denotes BCZY27 ICDD data, triangle symbol denotes LNICDD data.

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Journal of The Electrochemical Society, 161 (6) F717-F723 (2014) F719

25 30 35 40 45 50

Rela

�ve

Inte

nsity

(Arb

itrar

y U

nits

)

Angle (2θ Degrees)

Figure 3. XRD pattern of La2NiO4+δ powder after exposure to 15% steam inair at 700◦C. Symbol denotes LN ICDD data.

migration of lanthanum into the BCZY27 electrolyte exists, as shownin Figure 5. It is evident that the nano-particles are lanthanum nicke-late, and there appears to be no significant lanthanum migration intothe BCZY27. Figure 4(c) and 4(d) show a post-test cross-section. NoLN nano-catalyst growth appears during testing, which is a favorableresult as it indicates no significant change in catalyst active surfacearea. The LN catalyst appears to be stable at the high steam partialpressures and low oxygen partial pressures that it was exposed toduring the symmetric cell testing.

To determine the backbone porosity, five different sections of thebackbone were imaged and analyzed with EDAX Genesis software(edax.com). The average porosity was 41.9% with a standard devia-tion of 6.0. The amount of LN catalyst added to the backbone wascalculated using the pore volume and molarity of the infiltrate solu-tion. If it is assumed that the infiltration solution fills all pores, theweight of catalyst can be estimated at 13.3 mg over 10 infiltrations toa 7.3 wt% loading (183 mg of BCZY27 backbone).

Electrochemical characterization.—Data verification.— Kramers-Kronig tests show a deviation between the calculated and measureddata of much less than ±1% for both the LN cell and the controlcell. Figure 6 shows an EIS curve analyzed with the K-K test.In the case of the 600◦C, 21% O2, 3% steam measurement, themaximum error for the real and imaginary impedance is 0.2%. Thissuggests that there is no significant drift or noise in the measureddata.Fitting and analysis of EIS data.—Most of the data was fit to anequivalent circuit of R(RQ)(RQ)(RQ)(RQ), that is an ohmic resistancefollowed by a series combination of four parallel RQ circuits, where Qis the constant phase element (CPE). However, at 500◦C, some of theEIS spectra did not contain the required information to fit the last RQdue to a limitation in the low frequency of the measurement apparatus.Therefore, the data for the first two RQ elements are compared for alltemperatures (500, 600, and 700◦C), and the data for the last two RQelements are only compared for the 600 and 700◦C conditions. Theoverall polarization resistance is calculated by the sum of all resistiveelements. The polarization resistance is converted into an ASR bymultiplying by the active area of 1.23 cm2.

A typical fit of the EIS data using the Zsimpwin software is shownin Figure 6(a). A circuit consisting of five elements (a resistor and fourRQ elements) can be fit with many solutions, meaning there is not al-ways a unique solution. To ensure that data is being fit in a realisticmanor, equivalent circuit analysis must be performed and graphed inSchouler-type representation.16 This determines if the equivalent ca-pacitance is consistent as a function of reciprocal temperature, visuallyverifying that the same phenomena is represented for each equivalentcapacitance.

The calculated equivalent capacitances can be seen in Schouler-type representation in Figure 7. The consistent values for each capac-

5 m

a)

b)

100 nm

1 m

c)

d)

Figure 4. Pre-test SEM images of a) LN infiltrated BCZY27 backbone, b)LN nano-particles. Post-test images of c) infiltrated backbone, d) LN nano-particles.

itance as a function of temperature suggests that the same phenomenais being described at each temperature with a reliable and repeatablefit.

To understand the physical processes affected by the LN catalyst,it was imperative to conduct the same experiments on an identicalcell without the catalyst. This decouples backbone and catalyst pro-cesses. The control cell was only exposed to the experimental matrix at600◦C but at each of the 9 different gas compositions. Assuming that

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F720 Journal of The Electrochemical Society, 161 (6) F717-F723 (2014)

a) b)

BCZY27

Infiltrate

Figure 5. a) SEM image of EDS section, b) overlaid La element map.

a)

b)

Frequency (Hz)

Erro

r (Pe

rcen

t)

1%

-1%

0%

5.0 E3 1.0 E6

RRQ RQ RQ RQ

Figure 6. a) EIS data at 600◦C, 21% O2, 3% steam with R(RQ)(RQ)(RQ)(RQ) fit, b) K-K transform.

Figure 7. Schouler-type representation of equivalent capacitances as a func-tion of reciprocal temperature, 21% O2, 3% steam at 600◦C.

O2(g)

H+

O2-(ad)e-

H20(g)

OH-(ad)

H+

e-

H(ad)H2(g)

O2(g)H20(g)

Silver Current Collectora) b)e- from currentcollector

Figure 8. Diagram of the differing electrochemical mechanisms between a)infiltrated LN electrode and b) control electrode.

the silver current collector is inert to oxygen reduction reaction, thereactions can be visualized as in Figure 8. With the nano-catalyst in-filtrated into the backbone (Figure 8(a)), the reaction mechanism isa visually adapted version of the proposed mechanism by He et al..8

Electrons from the current collector are transported to the three-phaseboundary through the nano-catalyst to initiate electrochemistry. How-ever, with no catalyst in the control cell, a different reaction mechanismcan be hypothesized(Figure 8(b)). In this case, electrons cannot mi-grate to the electrolyte surface due to the lack of an electronic phase.Therefore, it is theorized that the proton must diffuse through theentire thickness of the backbone in order to reach electrons at the cur-rent collector, finally forming hydrogen. Without the presence of anoxygen-adsorbing catalyst, it is likely that the hydrogen-oxygen reac-tion takes place in the gas phase. The differences in the electrochemicalprocesses of these two cells were captured and analyzed by EIS.

Figure 9 compares two EIS spectra of the LN cell (a), and controlcell (b). Four RQ elements are apparent in the control cell, but it isfar more difficult to see how they manifest in the infiltrated LN cell.It is clear that the electrolyte conductivity is the same, as the ohmicresistance (Rs) value is nearly identical between the two samples.However, three of the other four RQ elements are far smaller in the LNcell, indicating that these are the elements involved in catalytic, chargetransfer, or electronic conduction events. The sum of the resistancesof the four RQ circuits is lower in the LN cell than the control cell,suggesting that the LN catalyst is active in reducing the polarizationresistance (Rp). The fourth RQ element remains nearly the same inboth cells.

a)

b)

Rs Rp

ControlLN

Figure 9. EIS spectra for LN-infiltrated and control cells at 600◦C, 21% O2,3% H2O: a) comparison, b) LN-electrode with Rs and Rp illustrated.

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Journal of The Electrochemical Society, 161 (6) F717-F723 (2014) F721

Table II. Calculated equivalent capacitance and relaxationfrequencies for the control and LN cells.

Calculated parameter Control electrode value LN electrode value

C1 2.3 × 10−8 F 4.6 × 10−8 Ff1 7.7 × 104 Hz 3.4 × 105 Hz

C2 1.8 × 10−5 F 1.3 × 10−6 Ff2 1.3 × 102 Hz 1.3 × 104 Hz

C3 1.5 × 10−2 F 3.2 × 10−1 Ff3 7.7 × 10−1 Hz 2.3 × 10−2 Hz

C4 6.9 × 10−1 F 4.2 × 100 Ff4 1.3 × 10−2 Hz 2.2 × 10−2 Hz

For the spectra shown in Figure 9, the equivalent capacitances andrelaxation frequencies are listed in Table II. Examining the equivalentcapacitances in the table, C1 likely describes the same phenomena be-tween the control electrode and the LN electrode. The resistance andrelaxation frequency values for C1 differ between the two electrodes,but this is easily explained. C1 is likely grain boundary resistance ofthe BCZY27 backbone, which would be present in both electrodesdue to the small grains of the backbone. However, the resistance isexpected to be smaller in the LN electrode, as the proton diffusionpath is shorter since electrons are carried to the protons by the elec-tronically conductive catalyst, as Figure 8 illustrates. C2 and C3 wereboth dramatically reduced in the LN cell as compared to the controlcell. This suggests C2 and C3 are catalytic events. Minimal resistivechange was observed for C4 between the LN cell and control cell.This result indicates that C4 is likely gas diffusion resistance due tothe microstructure of the backbone. The following conclusions canbe stated using the above analysis and equivalent capacitance valuesfrom literature:15-17

� C1 corresponds to grain boundary resistance in the BCZY27backbone;

� C2 corresponds to charge-transfer chemistry;� C3 corresponds to oxygen adsorption, desorption;� C4 corresponds to diffusion.

Electrochemical performance of lanthanum nickelate.—The contribu-tions to total polarization resistance are broken out into four separateASRs. Figure 10 shows the trend of the ASR of each contribution asa function of pO2. The 500◦C data was excluded from this analysisdue to the previously mentioned difficulty in resolving the fits for thedata at low frequency. The general trend is decreasing resistance withincreasing pO2 for all of the contributions. Also, grain-boundary andionic-transfer mechanisms have a larger temperature dependence thanthe dissociation/adsorption/diffusion mechanisms.

Although every elementary oxygen reaction listed by He et al.8

cannot be individually expressed in the EIS fit, Figure 11(a) illustratesthat the sum of all the oxygen adsorption/dissociation steps (labeledOX) has a measured reaction order of m = 0.17 at 600◦C. Similarly,the sum of all the charge transfer steps (labeled CT) has a measuredreaction order of m = 0.15 at 600◦C. It is also important to note that theoxygen adsorption/dissociation ASR is roughly double of the chargetransfer ASR, even though the reaction orders are very similar. Thissuggests that catalytic surface reactions are rate-limiting as opposedto electrochemical charge-transfer reactions.

The sum of all of these ASR values is equal to the total ASR, whichis shown in Figure 11(b) as a function of pO2. When compared to theindividual ASRs, the total ASR follows a more evident decreasingtrend with increasing pO2. This result is indicative of active oxygen-reduction-reaction catalysis of the electrode. It is important to notethat the contribution of the BCZY27 backbone grain-boundary resis-tance is significant in the total ASR. The grain-boundary resistanceis independent of LN catalyst, so minimization of grain-boundaryresistance through processing techniques is an area for improvement.

0

2

4

6

8

10

12

14

0 0.5 1

Gra

in B

ound

ary

ASR

(Ohm

s cm

2 )

pO2 (atm)

600 °C

700 °C

0

2

4

6

8

10

12

14

0 0.5 1

Char

ge T

rans

fer A

SR (O

hms c

m2 )

pO2 (atm)

600 °C

700 °C

0

5

10

15

20

25

30

0 0.5 1

Oxy

gen

Inte

rac�

on A

SR (O

hms c

m2 )

pO2 (atm)

600 °C

700 °C

0

5

10

15

20

25

30

35

0 0.5 1

Gas

Diff

usio

n AS

R (O

hms c

m2 )

pO2 (atm)

600 °C

700 °C

Figure 10. Comparison of ASRs as a function of pO2 for LN cell.

Although the stability of LN in steam was verified by XRD(Figure 2), significant degradation was observed during EIS char-acterization under varying pH2O. This degradation was not evidentafter varying only pO2. The trend of polarization resistance as a func-tion of steam partial pressure did not exhibit a decrease in resistancewith increasing steam as expected from estimated reaction orderspublished in literature.8 Instead, increasing ASRs as a function ofincreasing pH2O were observed. Figure 12 illustrates degradation of

Figure 11. a) Processes order of charge-transfer and oxygen adsorp-tion/dissociation at 600◦C, b) total polarization resistance of LN electrodeas a function of pO2.

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F722 Journal of The Electrochemical Society, 161 (6) F717-F723 (2014)

Figure 12. EIS spectra illustrating cell degradation before and after exposureto high steam, 600◦C, 21% O2, 3% steam.

200 m

a) b)

Electrolyte

Backbone

Silver Current Collector

Figure 13. Backscatter SEM images of a) Silver on control cell before steamexposure, b) silver on control cell after 15% steam exposure in air at 700◦C.

the LN symmetric cell due to testing in a high-steam environment.One EIS curve was taken at 3% steam before any steam increase, andthe other taken at 3% steam after the high-steam tests. The fact thatdegradation is evident agrees with the theory of steam degradation.However, previous XRD studies showed no degradation of the LN insteam, and post-mortem SEM verified no significant particle growthof the LN nano-particles.

To determine if current collection was problematic, a second con-trol cell was sectioned into two pieces. Silver paste was painted on thebackbone of each sample and sintered to 700◦C in air. After sintering,one sample was placed in 15% steam in air at 700◦C for five days.Both samples were mounted, polished, and examined using electronmicroscopy. Figure 13 shows the difference in microstructre betweenthe sample exposed to only ambient air (a), and the sample exposedto 15% steam at 700◦C (b). It is clear that the grain size of the silver islarger in the sample exposed to steam, but that can be explained by theadditional five days at temperature. However, there is a delaminationthat grows between the backbone and the silver current collector whenexposed to steam. It is believed that this growing delamination is thecause of the degradation in electrode performance under high-steamenvironments. Unfortunately, this degradation makes the high-steamresults untrustworthy, and they cannot be used to characterize infil-trated lanthanum nickelate in high-steam environments.

A list of all of the measured total polarization resistances as afunction of pO2 at 3% steam can be found in Table III. Again, it isevident that the polarization resistance decreases with both increasein temperature and increase in pO2.

Table III. Summary of measured ASR values as a function of pO2,3% steam.

pO2 500◦C ASR 600◦C ASR 700◦C ASR(atm) (�-cm2) (�-cm2) (�-cm2)

0.01 181.2 42.1 11.10.21 161.8 29.3 3.70.50 158.7 20.6 4.01.00 61.0 19.0 2.6

Figure 14. Electrolyte conductivity of BCZY27 as a function of pO2.

Quantification of electrolyte conductivity.—By measuring the Rs

value, the conductivity of the BCZY27 electrolyte can be determinedas a function of oxygen partial pressure. BCZY27 behavior as a func-tion of oxygen partial pressure can also be predicted by:

Oxygen incorporation reaction:

1/2O2(g) + V••O

⇀↽ O×O + 2h• [4]

We can re-write this equation in equilibrium constant form:

K = [O×O][h•]2

[O2(g)]1/2[V••O ]

[5]

Solving for the hole concentration gives:

[h•] = K1/2[O2(g)]1/4[V••

O ]1/2[O×O]−1/2 [6]

where it is evident that that hole concentration (and thus conductivity)is related to oxygen to the quarter power.

However, the measured conductivity of the BCZY27 cell doesnot follow this trend, as seen in Figure 14. This is for two reasons:1) competition for the oxygen-vacancy sites exists due to Wagnerhydration attempting to occupy oxygen vacancies (Eqn 7), and 2)other ohmic resistances, such as grain boundary resistance, lanthanumnickelate electronic resistance, and interfacial resistance decrease theconductivity. The additional resistances are not necessarily a functionof oxygen partial pressure, which reduces the overall partial pressuredependence.

Wagner hydration:

H2O(g) + V••O + O×

O⇀↽ 2OH•

O [7]

The theory that Wagner hydration and oxygen incorporation com-pete for oxygen vacancies is supported by the increase in the slope ofthe conductivity as a function of temperature in Figure 14. As temper-ature increases, dehydration of the OH•

O species occurs. This meansthat more oxygen vacancies can be occupied by oxygen incorporation,and the slope of the conductivity will increase toward the theoreticalquarter power calculated for pure oxygen environments.

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Journal of The Electrochemical Society, 161 (6) F717-F723 (2014) F723

Conclusions

Symmetric-cell electrolytes were fabricated using solid-state reac-tive sintering. Porous 40 μm thick BCZY27 backbones were sinteredto both sides of the electrolyte. Lanthanum nickelate (La2NiO4+δ) wassuccessfully introduced into the porous BCZY27 backbone by nitratesolution infiltration. Ten infiltrations of the 1 M nitrate solution re-sulted in a percolating catalyst loading of 7.3 wt%. XRD revealedno additional phases due to the low calcining temperatures used forthe LN infiltration. The infiltrated-backbone approach resulted in awell-adhered, phase-pure electrode.

EIS measurements were taken on the symmetric cell as a func-tion of pO2 and pH2O. The pH2O data was found untrustworthydue to the steam degradation of the silver current collectors. Therewas no significant degradation observed during the oxygen partialpressure study. As expected, the total polarization resistance reducedwith increasing temperature. More importantly, it was discoveredthat the overall polarization resistance reduced with increasing oxy-gen partial pressure. By fitting the data to a R(RQ)(RQ)(RQ)(RQ)circuit, the individual contributions of specific mechanisms wereidentified, including BCZY27 backbone grain-boundary resistance,charge-transfer, oxygen adsorption/dissociation, and gas diffusion. Atotal ASR of 3.7 �·cm2 was measured at 700◦C in 3% humidifiedair. This value is slightly higher than values from literature usingsingle-phase La2NiO4+δ electrodes.16,25 However, a significant contri-bution (∼25%) to the total ASR in the infiltrated cathode comes fromthe BCZY27 backbone grain-boundary resistance. By improving theBCZY27 backbone sintering, the total ASR of the cathode could besignificantly reduced.

Acknowledgments

This work has been funded by the Colorado School of MinesProtonic Capital Fund. Special acknowledgement to Dr. AnthonyManerbino and Dr. W. Grover Coors at CoorsTek, Inc. for techni-cal support.

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