Synthesis and Characterization of Sr[sub 2]MgMoO[sub 6−δ]

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Synthesis and Characterization of Sr2MgMoO6−�

An Anode Material for the Solid Oxide Fuel CellYun-Hui Huang, Ronald I. Dass, Jonathan C. Denyszyn, andJohn B. Goodenoughz

Texas Materials Institute, The University of Texas at Austin, Austin, Texas 78712, USA

The double-perovskite Sr2MgMoO6−� �SMMO� was investigated as an anode material of a solid oxide fuel cell. Via a syntheticmethod based on thermal decomposition of metal complexes with ethylenediaminetetraacetic acid as the complexant, phase-pureSMMO was readily obtained. Oxygen vacancies are introduced by reduction with 5% H2 at 800°C. With a 300 �m thickLa0.8Sr0.2Ga0.83Mg0.17O2.815 disk as the electrolyte and SrCo0.8Fe0.2O3−� as the cathode, the SMMO anode showed power densitiesof 0.84 W/cm2 in H2 and 0.44 W/cm2 in CH4 at 800°C. Moreover, it performed stably on power cycling and tolerated sulfur andmoisture well. Only 1% degradation in the output was observed in H2 containing 5 parts per million �ppm� H2S and 16%degradation in H2 containing 50 ppm H2S compared with the output in pure H2. Thermogravimetric analysis showed a drop inmass at around 750°C in the atmospheres of both air and 5% H2, indicative of the formation of oxygen vacancies. The meanthermal expansion coefficient was � = 12.7 � 10−6 K−1 at the operating temperatures. The conductivity strongly depended on theatmosphere, and the electronic activation energies were Ea = 0.084 eV in H2 and 0.126 eV in CH4. Our results show that SMMOis a potential anode material for operation with natural gas.© 2006 The Electrochemical Society. �DOI: 10.1149/1.2195882� All rights reserved.

Manuscript submitted December 21, 2005; revised manuscript received February 24, 2006. Available electronically May 4, 2006.

0013-4651/2006/153�7�/A1266/7/$20.00 © The Electrochemical Society

The solid oxide fuel cells �SOFCs� under development todaymostly operate on syngas, a mixture of H2 and CO, as the fuelbecause the best anode material is a cermet of NiO and a solidoxide-ion conductor. Reduction of the NiO in the H2 atmosphere atthe anode leaves a porous oxide-ion conductor with elemental nickelon the surface of the pores. The nickel is an excellent catalyst forbreaking the H2 bond, and it conducts the two electrons of thechemisorption reaction to the current collector; the two H+ ions arespilled over to the oxide-ion conductor to release H2O, and replace-ment O2− ions are fed from the cathode through the electrolyte to theoxide-ion conductor of the anode. However, nickel becomes fouledwith coke on the oxidation of a hydrocarbon fuel unless a largeamount of steam is added to the fuel,1,2 and sulfur impurities in thefuel react with nickel to form NiS, which requires a high-gradedesulfurization for the fuel.3 Therefore, use of a hydrocarbon fuelrequires a reformer that feeds syngas to the anode of the SOFC. Inorder to eliminate the reformer and to simplify the removal of sulfurfrom a hydrocarbon fuel, it is important as a first step toward the useof logistic fuels to identify an anode material for the SOFC thatprovides efficient operation with natural gas at 800°C. The anodematerial must be not only catalytically active for breaking the C–Hbond with subsequent release of CO or CO2 and H2O; it must alsobe able to conduct the electrons released to the current collector ofthe anode. For this purpose, a mixed oxide-ion-electron conductor�MIEC� would be ideal, although an alternative would be a catalyti-cally active oxide-ion conductor impregnated with CuO that is sub-sequently reduced to elemental copper for electronic conduction.

Early experiments by Gorte and his co-workers investigated rare-earth-doped ceria impregnated with CuO as the possible anode ma-terial for direct utilization of hydrocarbon fuels in SOFCs.4,5 Re-cently this material has been shown to be tolerant to 450 ppm H2S ata fixed potential of 0.65 V.6 These pioneering studies showed thefeasibility of such an approach, but the catalytic activity of thedoped ceria proved to be disappointing. Subsequently, Tao et al.7,8

reported that the oxygen-deficient perovskite system�La1−xSrx�0.9Cr0.5Mn0.5O3−� is an MIEC anode with comparable per-formance in H2 to Ni-based cermets and with catalytic activity forthe electro-oxidation of CH4 at high temperatures. However, it is notstable to sulfur impurities in the fuel;9 even when impregnated withCuO to improve the electronic conductivity in the reducing atmo-sphere at the anode, the long-term performance has provenunsatisfactory.10 Nevertheless, this study pioneered the use of an

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address. Redistribution subject to ECS terms129.10.52.112aded on 2014-04-06 to IP

oxygen-deficient, mixed-valent perovskite as an MIEC of potentialinterest for the anode material of an SOFC. These oxides also havea thermal-expansion coefficient that is compatible with that of thesolid oxide electrolyte. We have chosen to study the double-perovskite system Sr2Mg1−xMnxMoO6−� because it is oxygen defi-cient and stable in a reducing atmosphere.11 Moreover, the mixed-valent Mo�VI�/Mo�V� subarray provides electronic conductivitywith a large enough work function to accept electrons from a hydro-carbon. As an MIEC that can accept electrons while losing oxygen,it also promises to be catalytically active to the oxidation of H2 andhydrocarbons. The ability to lose oxygen while accepting electronsis realized because Mo�VI� and Mo�V� form molybdyl ions, whichmakes them stable in less than sixfold oxygen coordination, andboth Mg2+ and Mn2+ are stable in fourfold as well as sixfold oxygencoordination. We have found that Sr2MgMoO6−� gives the best per-formance, and we report here our structural, chemical, and transportcharacterizations of this material as well as its performance as ananode with a 300 �m thick La0.8Sr0.2Ga0.83Mg0.17O2.815 �LSGM�electrolyte and a SrCo0.8Fe0.2O3−� �SCF� cathode operating on H2and CH4 fuel with or without H2O or H2S. Our results show it is apotential anode material for operation on natural gas.

Experimental

Synthesis.— Sr2MgMoO6−� was prepared by a sol-geltechnique with ethylenediaminetetraacetic acid �EDTA� as thechelating agent and Sr�NO3�2, Mg�NO3�2·5.65H2O, and�NH4�6Mo7O24·4.00H2O as the starting materials. The water ofcrystallization of the precursor salts was determined by thermogravi-metric analysis �TGA� with a heating rate of 1°C/min in flowing air.Stoichiometric quantities of the salts were first dissolved in deion-ized water to give a clear solution while EDTA was separately dis-solved in aqueous ammonia; the molar ratio between the totalamount of metal ions and the amount of EDTA was 1:1.5. Theaqueous NH4-EDTA solution was then slowly added at room tem-perature to the stirring clear solution of metal cations; the pH of thissolution was adjusted to approximately 10 with the further additionof aqueous NH3. The solution was then heated to approximately150°C to allow the chelates to undergo polyesterification as well asto remove excess water. The resulting off-white-colored gel wascompletely dried in an oven to give a dark-brown, spongy gel thatwas then ground and slowly decomposed in air at 400°C for 24 h.The fine black powder produced was ground and annealed in air at800°C for 6 h. The product, a white powder with a tinge of green,

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was ground, pressed into half-inch diameter pellets 6–8 mm thick,and annealed at 1100°C in a flowing atmosphere of 5% H2/Ar at1100–1200°C for 24 h.

The electrolyte LSGM was prepared by conventional solid-statereaction. Stoichiometric La2O3, SrCO3, MgO, and Ga2O3 weremixed and ground well. La2O3 and MgO were dried in the furnaceat 980°C overnight before using. The mixed powder was pressedinto pellets and then calcined at 1000°C for 20 h. After coolingdown, the pellets were reground, pelletized, and calcined at 1200°Cfor 20 h in air. The thus-obtained precursor was ballmilled into veryfine powder, mixed carefully with 1 wt % polymer binder �PVB�,and then pelletized into disks with a diameter of 2 cm. The diskswere finally sintered at 1450°C for 20 h to achieve a pure perovskite

phase with a cubic symmetry of Pm3̄m space group �no. 221�. Thesintered ceramic disks were polished automatically on a machinewith a diamond wheel until thin and smooth. The thickness wascontrolled to 300 �m. The cathode material SCF was also preparedby conventional solid-state reaction. Stoichiometric SrCO3, Co3O4,and Fe2O3 were mixed, ground, and pelletized. First, the pelletswere calcined at 900°C for 10 h, then at 1000°C for 10 h, andfinally sintered at 1250°C for 10 h in air. It also showed a cubic

symmetry with space group of Pm3̄m �no. 221�. La0.4Ce0.6O2−�

�LDC�, a buffer layer material between the anode and the electro-lyte, was obtained via a sol-gel method based on the decompositionof EDTA complex gel. La2O3 and CeO2 were used as startingchemicals, and final sintering was carried out at 1450°C for 24 h inair. The phase was indexed to a cubic symmetry with space group

Fm3̄m �no. 225�.

Fabrication and testing of single fuel cells.— Single test fuelcells were fabricated by an electrolyte-supported technique. AnLSGM disk with a fixed thickness of 300 �m was used as the elec-trolyte. A thin LDC buffer layer was used between the anode and theelectrolyte to provide isoactivity of lanthanum at the buffer-electrolyte interface;12 the fluorite structure of the buffer layer pre-vents interdiffusion of ionic species between the perovskite anodeand electrolyte. The thin buffer layer has proven effective in numer-ous studies with the LSGM electrolyte.12-14 Without LDC as thebuffer layer, the power density of the cell drops to almost half thevalue of that with the buffer layer. LDC, SMMO, and SCF weremade into inks with a binder, V-006 �Heraeus�. LDC ink was screenprinted onto one side of the LSGM disk followed by firing at1300°C in air for 1 h. SMMO was subsequently coated on the LDClayer and baked at 1275°C in air for 1 h. SCF was finally screenprinted on the other side of the LSGM disk and fired at 1100°C inair for 1 h. The electrode area of the cell was 0.24 cm2. Becausesilver and gold pastes melt at around 800°C, which remarkablyreduces the effective electrode area, we used Pt mesh with a smallamount of Pt paste as a current collector at both anode and cathodesides for ensuring contact. Reference electrodes of the same materi-als as the working electrodes were used to monitor the overpoten-tials of the cathode and anode in the cell configuration. In order toavoid potential gradients along the electrolyte surface, the referenceelectrodes were placed greater than eight electrolyte thicknessesaway from the working electrode.15 A double-layer sealing designwas applied to the single cells. A Pyrex glass ring was placed be-tween the cell and a supporting Al2O3 tube for the internal sealing.A slurry, made by mixing Duco cement and Pyrex glass powder withacetone, was pasted on the outside of the Pyrex glass ring to serve asthe external sealing. The applied glasses were softened at about750°C to seal the whole cell.

Characterization and performance test.— The phases of thesamples were identified with a Philips X-pert powder X-ray diffrac-tometer �XRD� and Cu K� radiation. A Perkin-Elmer series 7 ther-mogravimetric analyzer �TGA� was used to monitor the oxygen de-ficiency of SMMO under operating conditions. Micrographs weretaken by a scanning electron microscope �SEM, Hitachi: S4500�.


The conductivity was measured by a standard dc four-probe methodwith our own setup. Pt wire and Pt paste were used to make the fourprobes. The samples were pressed and polished into rectangular barsand then sintered in 5% H2/Ar at 1150°C for 24 h. Before measure-ment, the sample was again reduced in 5% H2/Ar at 800°C for 20 hto ensure formation of oxygen vacancies. As a function of oxygenpartial pressure pO2, the conductivity was measured isothermally at800°C. The measurements started from pure H2 and then thesamples were slowly allowed to reoxidize with a slow air leak tochange pO2. Thermox CG1000 oxygen analyzer �Ametek� was usedto monitor the oxygen partial pressure.

For the performance test, the assembled cells were placed in thehot zone of a vertical furnace with air directly supplied to the cath-ode surface and the fuel to the anode surface at a flow rate of30 mL/min. Before testing, the cells were exposed to 5% H2/Ar for20 h at 800°C to reduce SMMO and then purged with fuel gas for2 h. The performance measurements were typically carried out inthe temperature range from 650 to 850°C, and a constant potentialwas provided by an EG&G potentiostat-galvanostat model 273 run-ning on a homemade LabView program. During a typical measure-ment, the cell voltage was varied from open circuit voltage �OCV�,which is around 1.2 V in H2 and 1.0 V in CH4, to 0.4 V and back toOCV in a total of 30 steps and holding 10 s at each step.

Results and Discussion

Figure 1 displays XRD patterns for Sr2MgMoO6−� samples ob-tained under different sintering conditions. The pure double-perovskite phase is readily achieved. Even for a sample sintered at900°C in air, only a tiny amount of impurity was observed. Cellperformance tests show no obvious difference in the power densityfor the cells with the SMMO samples sintered at different tempera-tures from 1100 to 1200°C in 5%H2/Ar. Therefore, in the presentwork, we investigated the properties of samples sintered at 1200°Cfor 24 h. From the XRD patterns, SMMO is monoclinic �spacegroup P21/n, b-axis unique� characteristic of a well-ordered doubleperovskite with lattice parameters a = 5.5920�2�, b = 7.8839�3�,c = 5.5871�7�, and � = 89.693�5�°.

After reoxidation of Sr2MgMoO6−� in air in a TGA, a weightdecrease was observed on heating at about 750°C in both air and 5%H2/Ar, as displayed in Fig. 2. This change corresponds to an escapeof oxygen from the lattice, which gives rise to oxygen vacancies.The thermal expansion behavior of SMMO in air is shown in Fig. 3.

Figure 1. XRD patterns for Sr2MgMoO6−� samples obtained at differentsintering temperatures for 24 h: �a� 900°C in air, �b� 1100°C in 5%H2, �c�1150°C in 5%H2, �d� 1200°C in 5%H2, and �e� anode film after testing inH2, H2/H2S, and CH4.




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Two different linear relationships between the sample heightand temperature were observed. The two linear segments gavethermal expansion coefficients � = 12.7 � 10−6 K−1 �691 K � T� 1074 K� and � = 11.7 � 10−6 K−1 �382 K � T � 633 K�,which are comparable to that of YSZ �10.3 � 10−6 K−1� and that ofPt �9.0 � 10−6 K−1�.

Figure 4 presents Arrhenius plots of the conductivity � ofSr2MgMoO6−� in 5%H2/Ar, H2, H2/5 ppm H2S �briefly, H2/5H2S�and CH4 after reduction in 5% H2/Ar at 800°C for 20 h. In allatmospheres, the sample exhibits polaronic conduction. The � valuedepends strongly on the reduced state of the sample. The sampleas-prepared in 5% H2/Ar without further reduction shows a verylow conductivity. However, after it is further reduced in 5% H2/Arfor 20 h, � increases by almost 3 orders of magnitude. As indicatedby the TGA result, reduction in 5% H2/Ar leads to the formation ofoxygen vacancies in the SMMO structure. Each oxygen vacancyreduces 2 Mo�VI� to Mo�V�. The enhanced conductivity can beascribed to a mixed-valent Mo�VI�/Mo�V� couple that provides agood electronic conduction. At 800°C, the � value is 4 S cm−1 in5% H2/Ar, but close to 10 S cm−1 in H2. Figure 5 shows the elec-tronic conductivity measured in H2 with leaking air at 800°C as afunction of oxygen partial pressure pO2. With increasing pO2, theconductivity decreases, indicative of an n-type conductivity as the

Figure 2. TGA run of Sr2MgMoO6−� recorded in air and 5%H2/Ar with aheating rate of 1°C/min.

Figure 3. Thermal expansion behavior of Sr2MgMoO6−� in air �the height ofthe sample cylinder was 5.1337 mm at 299 K�.


dominant electronic mechanism. However, � is not sensitive to lowpO2 values. The temperature dependence of the conductivity can bedescribed by a small-polaron hopping mechanism

� = �A/T� exp�− Ea/kT� �1�

where Ea = �Hm + ��Ht/2� is the sum of the motional enthalpy�Hm of the polarons and the enthalpy �Ht to free a Mo�V� from theoxygen vacancy that creates it. The value of Ea can be obtained fromthe slopes of the Arrhenius plots of ln��T� vs 1/T. In the interval400–800°C, the values of Ea are 0.197 ± 0.002, 0.084 ± 0.001,0.103 ± 0.003, 0.096 ± 0.001, and 0.126 ± 0.002 eV in 5% H2/Ar,H2, 97% H2-3% H2O, H2/5H2S and CH4, respectively. Only a slightdecrease in conductivity was observed with 3% H2O in the fuel. Forexample, at 800°C, the values of � are 8.6 and 8.5 S cm−1 in H2 and97% H2-3% H2O, respectively. See Table I.

For the single cell with SMMO as the anode, clear LSGM/SCF,LSGM/LDC, and LDC/SMMO interfaces are distinguished �see Fig.6� after testing in H2, H2/5H2S, and CH4. From the SEM images,the thicknesses of the LDC, SMMO, and SCF layers are estimatedto be 30, 100, and 20 �m, respectively; the grain and particle sizes

Figure 4. Temperature dependence of the electronic conductivity ofSr2MgMoO6−� in 5% H2/Ar, H2, 97%H2–3% H2O, H2/5H2S, and CH4 afterreduction in 5%H2/Ar at 800°C for 20 h.

Figure 5. Electronic conductivity measured at 800°C as a function of oxy-gen partial pressure for Sr MgMoO .
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for LDC, SMMO, and SCF are distributed between 1 and 2 �m.Furthermore, the anode SMMO exhibits a fine uniform microstruc-ture with an estimated 30% porosity.

Figure 7 depicts cell voltage and power density as a function ofcurrent density at different temperatures for the single fuel cell withSr2MgMoO6−� anode in dry and wet H2, H2/5H2S, and CH4. In dryH2 �Fig. 7a�, the maximum power density Pmax reaches 0.84 W/cm2

at 800°C, 0.65 W/cm2 at 750°C, and 0.45 W/cm2 at 700°C, whichshows an excellent performance at an intermediate operating tem-perature. Even in wet H2 �97% H2-3% H2O�, Pmax is still as high as0.81 W/cm2 at 800°C. Only a slight drop in Pmax is observed, indi-cating a high tolerance to moisture. In dry H2/5H2S �Fig. 7b�, Pmaxis 0.83 W/cm2 at 800°C and 0.64 W/cm2 at 750°C, whereas in wetH2/5H2S, Pmax is 0.80 W/cm2 at 800°C. The values of Pmax are ashigh as 0.44 and 0.35 W/cm2 in dry and wet methane at 800°C,respectively. Such a high power density in dry methane indicatesthat Sr2MgMoO6−� has great potential to be directly used for meth-ane oxidation without steam. In most SOFC cells with natural gaslike CH4 as a fuel, CH4 reforms with steam on the anode intosynthesis gas

CH4 + H2O = CO + 3H2 �2�followed by electrochemical oxidation of hydrogen at the sameanode

H2 + O2− = H2O + 2e− �3�

For the Sr2MgMoO6−� anode, a very high power density was ob-tained in dry CH4. With steam, the performance became lower.Steam is not necessary for this anode. Apparently CH4 can be con-

Table I. Values of the activation energy Ea of electronic conditionfor 400–800°C and the electronic conductivity � at 800°C in dif-ferent atmospheres.

Atmosphere Ea �eV� � �S/cm�

5% H2/Ar 0.197 ± 0.002 4.26H2 0.084 ± 0.001 8.6097% H2 − 3% H2O 0.103 ± 0.003 8.50H2/H2S 0.096 ± 0.001 8.54CH4 0.126 ± 0.002 8.5397% CH4 − 3% H2O 0.149 ± 0.001 6.69

Figure 6. SEM images for the cell after testing in H2, H2/H2S, and CH4 at650–800°C: �a� cross section of the cell configuration, �b� section betweenSCF and LSGM, �c� section between LSGM and LDC buffer layer, and �d�the surface of the Sr MgMoO anode.
2 6−�

verted by direct electrochemical oxidation on the anode providedoxide ions are transferred through the electrolyte

CH4 + 4 O2− = CO2 + 2H2O + 8e− �4�

With an LSGM thickness of 300 �m and SMMO of 100 �m, wehad a Pmax = 0.84 W/cm2 in H2 at 800°C. Changing the LSGMthickness to 500 �m while keeping SMMO 100 �m thick, loweredP to 0.77 W/cm2. Changing the SMMO thickness to 50 or

Figure 7. SMMO/LDC/LSGM/SCF cell voltage and power density as afunction of current density in different fuels: dry and wet H2, H2/5H2S, andCH4. The open symbols represent the cell voltages while the closed symbolsrepresent the power densities.

max





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150 �m while keeping the LSGM thickness at 300 �m gave a Pmaxless than 0.6 W/cm2 in H2 at 800°C. Therefore, the optimum anodethickness was about 100 �m, and the cell performance improves asthe thickness of the LSGM is reduced.

Figure 8 shows anode and cathode overpotentials a and c asfunctions of current density in H2, H2/5H2S, H2/50 ppm H2S�briefly, H2/50H2S�, and CH4. All the c show good linearity withcurrent density, indicating that oxygen penetration through the cath-ode is not rate-limiting. The a in H2 is linear with the currentdensity, whereas those in H2/5H2S and H2/50H2S are almost linearexcept at a low current density, which demonstrates that the exis-tence of a small amount of H2S has little obvious influence on theanode reaction. The nonlinear dependence of a on the current den-

Figure 8. Overpotentials �a and c� vs current density for a cell operatingat 800°C in H2, H2/5H2S, H2/50H2S, and CH4.

Figure 9. The maximum power density Pmax at 800°C vs cycle number indry and wet H , H /H S, and CH .
2 2 2 4

sity is due to the surface-reaction kinetics on the anode. However,the value of a is still lower than that of c in the high currentdensity range in CH4, which means that SMMO is promising fordirect oxidation of CH4 in SOFCs.

We ran the cell for repeated power cycles from OCV to 0.4 Vand back to OCV to check the stability of the cell performance.Figure 9 is a measure of Pmax of the single cell at 800°C againstcycle number in different fuels. In order to examine the tolerance toa higher concentration of sulfur, H2/50H2S was also used as fuel. Indry fuels, Pmax fades by 3.5, 2.5, 4.8, and 16% over 50 cycles in H2,H2/5H2S, H2/50H2S, and CH4, respectively. In wet fuels, the corre-sponding fades in Pmax over 50 cycles are 6.3, 5.9, 2.6, and 10%. Itis thus concluded that SMMO performs stably over many powercycles even in fuels containing a small amount of sulfur and mois-ture. Furthermore, the Pmax curves vs cycle number in H2 andH2/5H2S essentially overlapped, indicative of an excellent toleranceto 5 ppm H2S. Although 16% degradation in the output inH2/50H2S was observed compared to that in pure H2, the output wasstable over 50 cycles. The SMMO anode exhibits a much highertolerance to sulfur than the extensively used Ni-based anodes thatare poisoned by H2S even at levels as low as 0.05 ppm.3 We at-tribute this to a strong resistance of the Mg2+ ions to sulfideformation.

Considering that both Pt and LDC are catalytically active foroxidation of methane, we examined with energy dispersive spectros-copy �EDS� the elemental distribution along the cross-sectional lay-ers of the Sr2MgMoO6−�/LDC/LSGM/SCF cell with Pt mesh and Ptpaste as current collector taken after cycling in dry and wet H2,H2/H2S, and CH4 at 800°C for a total of 10 days. Six typical areaswith a grid size of 1 � 1 �m were chosen for EDS analysis, whichare sketched as labels of A–F in Fig. 6. A was located in the SMMOlayer close to the anode surface, B in the SMMO layer close toLDC, C in the LDC layer close to SMMO, D in the LDC layer closeto LSGM, E in the LSGM layer close to LDC, and F in the deep

Figure 10. EDS for the Sr2MgMoO6−� anode taken after cycling in H2,H2/H2S, and CH4.





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LSGM layer. EDS spectra for A–F are shown in Fig. 10, and theatom percentages of all the detected elements are listed in Table II.The amount of O was much higher than that expected due to a largeamount of O2− that reacted in these areas. For areas A and B, theatom ratio of observable Sr, Mg, and Mo obtained from EDS wasclose to 2:1:1 of the original sample. The phase of the anode filmafter testing in H2, H2/H2S, and CH4 checked by XRD also indi-cated no phase change and no impurity, as shown in Fig. 1e. Theseobservations demonstrate that the Sr2MgMoO6−� anode performsextremely stably in the reducing atmospheres. No carbon was de-tected. Only tiny amounts of S and Pt were found in area A that wasclose to the anode surface. We emphasize that no Pt appeared any-where in the LDC layer and no Pt was found in the deep SMMOlayer. Therefore, it is impossible for Pt and LDC to catalyze theoxidation of CH4 together. Both C and D areas contained someSMMO and LSGM, and E contained some LDC, which illustratesthat SMMO and LSGM phases interdiffuse in the LDC layer. How-ever, no LDC was observed in area B, indicating that the LDC layerprotected the SMMO layer well. The main purpose of the LDC layeris to prevent SMMO and LSGM from interdiffusion.

Furthermore, we directly used LDC without SMMO as the anodeand Pt mesh and Pt paste as current collector. The power density andcell voltage at 800°C in different fuels are shown in Fig. 11. Thevalues of Pmax were 0.53 and 0.49 W/cm2 in dry and wet H2, re-spectively, much lower than those of the cell with Sr2MgMoO6−� asthe anode. In methane, the output became even lower: 27 mW/cm2

in dry CH4 and 52 mW/cm2 in wet CH4. LDC exhibited some cata-lytic effect on methane reforming with the existence of steam, but it

Table II. Elemental analysis by EDS for theSr2MMoO6−�/LDC/LSGM/SCF cell with Pt mesh and Pt paste ascurrent collector after operating in dry and wet H2, H2/H2S, andCH4 at 800°C for a total of 10 days.

Atom % A B C D E F

O 76.83 78.02 76.73 73.23 74.1 77.95Sr 11.76 11.43 4.99 3.33 2.95 2.67

Mg 5.56 5.42 2.64 1.5 / 1.89Mo 5.41 5.13 1.44 0.53 / /La /a / 5.77 10.08 10.88 8.5Ce / / 7.18 4.21 4.30 /Ga / / 0.72 7.12 7.78 8.99S 0.36 / / / / /Pt 0.09 / / / / /

a Means zero.

Figure 11. Power density and cell voltage at 800°C for LDC/LSGM/SCFcell without double-perovskite Sr MgMoO in different fuels.
2 6−�

worked poorly. For direct oxidation of dry methane, the performancewas even poorer. Our Sr2MgMoO6−� anode showed a Pmax of0.44 W/cm2 in dry methane without any steam, which indicates thatSr2MgMoO6−� plays an important role for direct oxidation of CH4.

To further check the catalytic effect of Pt paste, we buried aPt-mesh current collector into the Sr2MgMoO6−� anode before an-nealing at 1100°C for 1 h. No Pt paste was used. The performanceis displayed in Fig. 12. The Pmax values were 0.92 and 0.71 W/cm2

in H2 at 850 and 800°C, respectively, about 10% lower than those ofthe corresponding cell with Pt paste. In dry methane, Pmax alsoreached as high as 0.34 W/cm2 at 800°C. In wet methane, Pmax was0.28 W/cm2 at 800°C. The conductivity of our materials is good. Ptpaste can enhance the conductivity of the anode, especially the lat-eral conductance; however, the cell with SMMO as the anode stillworked well without Pt paste. Therefore, the effect of Pt paste on thecell performance in our case is much less than previously reported.16

Conclusions

The double-perovskite oxide Sr2MgMoO6−� provides excellent,stable performance as the anode of an SOFC operating on H2 as fueland a promising performance with CH4 as fuel. Moreover, it isresistant to poisoning by sulfur impurities to 50 ppm H2S in the fuel.In the reducing atmosphere at the anode, it is a mixed oxide-ion–electronic conductor with electronic conduction on the Mo�VI�/Mo�V� redox couple. In 5% H2/Ar atmosphere, oxygen vacanciesare introduced above 750°C. The ability to lose oxygen while re-ceiving electrons from the fuel makes it catalytically active. Inves-tigations on the effects of the La-doped ceria and Pt paste evidentlyindicate that the good performance is attributed to the double-perovskite anode materials.

Acknowledgment

We thank the Robert A. Welch Foundation, Houston, TX, forsupport of this work.

University of Texas at Austin assisted in meeting the publication costs ofthis article.

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Figure 12. Power density and cell voltage at 800°C forSr2MgMoO6−�/LDC/LSGM/SCF cell with a buried Pt mesh in theSr2MgMoO6−� anode in dry H2 and CH4.





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Synthesis and Characterization of Sr[sub 2]MgMoO[sub 6−δ]