Catalysis of the hydrogen oxidation reactions by Sr-doped LaMn1 − yCryO3 ± δ oxides

Solid State Ionics 203 (2011) 69–79

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Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i

Catalysis of the hydrogen oxidation reactions by Sr-doped LaMn1 − yCryO3 ± δ oxides L. Deleebeeck, V. Birss ⁎ Department of Chemistry, University of Calgary, 2500 University Drive N.W. Calgary Alberta Canada T2N 1N4

a r t i c l e

i n f o

Article history: Received 14 February 2011 Received in revised form 14 July 2011 Accepted 26 July 2011 Available online 13 October 2011 Keywords: Solid oxide fuel cells Anode La1 − xSrxMn1 − yCryO3 ± δ LSCM Catalytic activity Chemical stability Temperature programmed reduction Electrochemical impedance spectroscopy Cr content Sr content

a b s t r a c t Sr-doped and Sr-free La1 − xSrxMn1 − yCryO3 ± δ (LSMC, x(Sr) = 0–0.2, y(Cr) = 0.4–0.6) perovskite-type oxides were synthesized and evaluated as single phase anodes for use in intermediate temperature solid oxide fuel cell applications. Their thermo-chemical and chemical stabilities were investigated in hydrogen at high temperatures and correlated with their oxygen non-stoichiometry (3 ± δ), determined by permanganate titration. The catalytic activity towards hydrogen oxidation was examined as a function of oxide sintering time, operating temperature, and the Sr and Cr contents, using a Pt mesh current collector. While all of the perovskite oxides studied here showed some irreversible performance degradation with time under both open circuit and anodically polarized conditions, La0.9Sr0.1Mn0.6Cr0.4O3.03 (LSMC9164), sintered at 1200 °C for 10 h, was found to be the most catalytically active and also the most stable. © 2011 Published by Elsevier B.V.

1. Introduction Solid oxide fuel cells (SOFCs) are devices that produce electricity at a higher efficiency than do conventional gas turbines or other combustion-based technologies. When both heat and power are extracted from an SOFC, efficiencies of N90% have been reported [1]. Additionally, since air and fuel are never combined directly in an SOFC, these devices do not produce pollutants, such as NOx, SOx, and particulates, responsible for poor air quality, especially in urban environments. An SOFC consists of an anode and cathode, separated by an electrolyte layer, which allows the transfer of oxygen anions (O 2−), formed at the cathode, and prevents the exchange of electrons. Ionically (O 2−) conducting solid oxide electrolytes, such as yttriastabilized zirconia (YSZ), give these fuel cells their name. SOFCs function at elevated temperatures (600–1000 °C) in order to enhance the conductivity of the electrolyte layer. The cathode, where the oxygen (from air) is reduced, is typically a composite mixture of an electronically conductive metal oxide, such as La1 − xSrxMnO3 ± δ (LSM), and the ionically conducting electrolyte, e.g., YSZ, or more recently, mixed ionic and electronic conducting oxides (MIEC), such as (La,Sr)(Co,Fe)O3 − δ (LSCF) oxide [1].

⁎ Corresponding author. Tel.: + 1 403 220 6432; fax: + 1 403 289 9488. E-mail address: [email protected] (V. Birss). 0167-2738/$ – see front matter © 2011 Published by Elsevier B.V. doi:10.1016/j.ssi.2011.07.017

The anode is typically composed of a cermet, which is a mixture of YSZ ceramic and Ni metal. Ni provides both the electronic conductivity and the reaction sites for catalyzing the oxidation of the fuel (H2, CH4, or other hydrocarbons) at the anode [1]. However, Ni also has several disadvantages, such as poor oxidation/reduction (redox) cycling tolerance [2], a tendency to coke in hydrocarbon fuels without the addition of large amounts of steam [1], and irreversible poisoning by sulfur, even in low ppm H2S environments [3]. Due to these drawbacks, alternate anode materials are currently under investigation, including Cu–CeO2 cermets [1] and MIEC metal oxides, such as doped perovskite oxides ((A, A’)(B, B’)O3 ± δ [4]). La1 − xSrxMn1 − yCryO3 ± δ (LSMC) perovskite-type oxides are known to have reasonable catalytic activity towards the oxidation of hydrogen [4] and methane [5], along with good coke tolerance [6]. Under high temperature, reducing conditions (pO2 ~ 10 − 20 atm), typical of the SOFC anode environment, LSMC materials are MIECs [7], show good chemical stability when in contact with standard electrolyte and interconnect materials, have a matching thermal expansion coefficient (TEC) with YSZ [8], and undergo only a slight volume change of ~ 0.2% upon redox cycling [8]. While LSMC materials have been shown to coke under open circuit conditions [9], they are reported to be coke-resistant and capable of catalyzing the complete oxidation of CH4 when polarized [6,9]. In previous work, we have shown that Sr-free LaMn0.5Cr0.5O3-gadolina doped ceria (GDC) composite anodes also remain coke-free when operated under polarized conditions [10].

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LSMC materials have been reported to be unstable in environments containing high concentrations of H2S, e.g., 0.5% H2S in CH4 [11] and 10% H2S in H2 [12]. In both cases, the secondary phases formed at high temperature in the presence of H2S were MnS, α-MnOS and La2O2S [11,12]. However, La0.8Sr0.2Mn0.2Cr0.8O3–YSZ anodes experience only a 10 mV loss at 500 mA cm − 2 in 10 ppm H2S [13], suggestive of reasonable sulfur tolerance at low H2S concentrations. In our previous work [10], we examined the thermo-chemical stability, electronic conductivity, and catalytic activity of Sr-free (x(Sr) = 0) and x(Sr) = 0.2 La1 − xSrxMn0.5Cr0.5O3 under reducing conditions. We reported that the Sr-free oxide is more stable (less reducible), has the higher electronic conductivity, and possesses better electrochemical activity than the x(Sr) = 0.2 oxide when operating at 700 and 800 °C in humidified H2. In the present work, the effect of a broader range of the Sr and Cr content of the La1 − xSrxMn1 − yCryO3 ± δ oxides (x(Sr) = 0–0.2, y(Cr) = 0.4–0.6) on the physico-chemical properties of single phase LSMC oxides, as well as on their H2 oxidation activity (when using a Pt current collector), has been investigated. Composite anodes, containing both LSMC and an electrolyte phase, were not studied here, as our primary objective was not to achieve maximum performance, but rather to employ conditions that would more clearly elucidate the differences in the LSMC oxide characteristics as a function of Sr and Cr content. We confirm that LMC55 (x(Sr) = 0, y(Cr) = 0.5) is more catalytically active than LSMC8255 (x(Sr) = 0.2, y(Cr) = 0.5) [10], and demonstrate that LSMC9164 (x(Sr) = 0.1, y(Cr) = 0.4), sintered at 1200 °C for 10 h, is the most stable and catalytically active oxide of all of the materials investigated here at both 700 and 800 °C in humidified H2. 2. Experimental procedure 2.1. Material synthesis La1 − xSrxMn1 − yCryO3 ± δ (x(Sr) = 0, 0.1, 0.2 and y(Cr) = 0.4, 0.5, 0.6) oxides were synthesized by co-precipitation, as described previously [10,14]. Briefly, La(NO3)3•xH2O (Sigma Aldrich, 99.9%), Sr (NO3)2 (BDH, 99.0%), Mn(NO3)2•xH2O (Alfa Aesar, 99.98%), and Cr (NO3)3•9H2O (Alfa Aesar, 99.99%) were combined in exact molar ratios, determined by the desired perovskite composition, to achieve a final perovskite oxide weight of 30 g. The metal nitrates were dissolved in water and then NH4OH was added at a LSMC:NH4OH ratio of 1:5. This caused a gel to form, which was then heated on a hot plate to remove excess water, followed by heating (covered) at 310 °C for 1 h to remove all remaining traces of water. The resulting solid was ground in a mortar and pestle, calcined for 12 h at 1260 °C in air, and then the final product was again ground in a mortar and pestle for at least 20 min. 2.2. Composition verification In order to precisely determine the number of waters of hydration (x H2O) of all of the metal nitrates, known masses of the salts were heated at N100 °C for several hours, resulting in their dehydration without allowing thermal decomposition, and were then weighed again. For La(NO3)3•xH2O, x was found to be 5.6, while x for Sr(NO3)2 and Cr(NO3)3 was confirmed to be 0 and 9, respectively. However, this method was ineffective for Mn(NO3)2•xH2O, due to its thermal decomposition into multiple Mn oxides of uncertain composition. Therefore, the degree of hydration of Mn nitrate was determined by Inductively Coupled Plasma Atomic Emission Spectroscopy (ICPAES) using a Thermo Jarrell Ash AtomScan 16 (Department of Chemistry, University of Calgary). A known mass of Mn(NO3)2•xH2O was dissolved in water and then diluted, with water to 5 ppm Mn. The Mn signal was compared to that from a known Mn AAS/ICP standard (1000 ppm, Alfa Aesar), with x found to be ~ 4. However, due to the

challenges of this analysis, this is only an approximate value. Additionally, the metallic content of the LSMC oxides was verified by ICP-AES analysis. 10 mg of sample was refluxed in aqua regia (1:3 HNO3:HCl) for 24 h and then diluted with water until each element was calculated to have a lower than 5 ppm concentration. ICP-AES was conducted using known La, Sr, Cr and Mn standards (1000 ppm, Alfa Aesar) diluted to 5 ppm. Attempts were also made to determine the total Mn content and the average Mn oxidation state of each LSMC (y(Cr) = 0.4) oxide by permanganate titration [15,16]. The oxygen non-stoichiometry of these oxides was calculated using the mean Mn oxidation state and the La, Sr, Mn and Cr molar stoichiometries, determined by ICP-AES analysis. 2.3. Physical characterization pXRD experiments were carried out using a Rigaku Multiflex XRay Diffractometer (Department of Geosciences, University of Calgary), operated at 40 kV and 20 mA using a Cu target material. The pXRD data were acquired between 10 and 80 2-theta (2θ) values at a scan width of 0.01° and a scan rate of 0.5 2θ min − 1. Phase purity was determined by pXRD analysis of the as-synthesized compounds, after calcining at 1260 °C for 12 h in air. The chemical stability of the oxides was also determined by pXRD after exposure to air for 10 h at 1200 °C, thus simulating electrochemical half-cell sintering conditions, and after reduction at 800 °C in H2 (3% H2O) for 5 days, simulating the electrochemical testing conditions. The reducibility of the oxides, defined as the % mass loss at a given temperature, was determined by temperature programmed reduction (TPR) between 150 and 1100 °C in 20% H2:80% He at a ramp rate of 1 °C min− 1. TPR experiments were carried out on a Setaram TAG16 Thermo-Gravimetric Analyzer (TGA), equipped with a dual chamber balance, using Pt crucibles. Each experiment was post factum corrected for the weight of the empty crucibles. The phase formed following TPR was then analyzed by pXRD, with the composition reproduced by exposing each LSMC oxide, for 24 h, to 1100 °C in 100% H2. Finally, these LSMC materials were re-sintered at 1260 °C for 12 h in air, and then their structures were examined by pXRD to determine the reversibility of phase formation under extreme reducing conditions [14]. Cross-sectional images were obtained using Scanning Electron Microscopy (SEM, Phillips/FEI X23 ESEM (Health Sciences, University of Calgary)), operated in vacuum mode at an accelerating voltage of 20 kV. Samples were prepared by embedding fractured electrolytesupported half-cells, previously used for electrochemical evaluation, in epoxy cold mounting resin, hardened with triethylenetetramine (Electron Microscopy Science, Hatfield, PA). 2.4. Electrochemical measurements Commercial 8 mol% yttria-stabilized zirconia (8YSZ, Tosoh) powder was pressed into 1 mm thick electrolyte disks [14]. The LSMC materials were combined with α-terpineol and ethyl cellulose to form a paste, which was screen-printed (300 mesh) over a symmetrical, circular area (~0.4 cm 2) on either side of the sintered YSZ disks (Fig. 1). The LSMC-coated YSZ disks were then sintered for 2 to 10 h at 1200 °C in air. A Pt paste (Fritted Pt conductor paste, Heraeus Inc., Germany) reference electrode (RE) was then painted on the counter electrode (CE) side of the YSZ disk over a small, but variable, surface area and at a distance of 5 mm from the CE (Fig. 1). The RE was sintered in-situ during a ramp (4 °C min − 1) to the operating temperature (700 or 800 °C) in air in a horizontal tube furnace (Thermo Electron Corp.). The 3-electrode cell, sintered prior to application of current collectors, was press contacted using 52 mesh Pt mesh current collectors, each spot-welded to a Pt lead wire. The planar half-cell and the current collectors were then loaded into a half-cell holder, as

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Fig. 1. Schematic of symmetrical electrolyte-supported cell used for 3 electrode half-cell studies. Working electrode (WE) and counter electrode (CE) were both screen-printed LSMC anode materials (area ~ 0.4 cm2), while the reference electrode (RE) was painted Pt paste, all press-contacted to Pt mesh current collectors connected to Pt wires.

described elsewhere [15]. Electrochemical impedance spectroscopy (EIS), potentiostatic, and cyclic voltammetry (CV) studies were performed at 700 and 800 °C in H2 (3% H2O) using a 1287 Solartron Instruments Electrochemical Interface Potentiostat and a 1255 Solartron Instruments Frequency Analyzer. EIS was carried out at the open circuit potential (OCP) at frequencies from 10 5 to 0.005 Hz and back again, to ensure overlap, with an AC amplitude of 10 mV (rms). Potentiostatic studies were performed at + 0.6 V (vs. Pt RE) for 2 to 256 min, with one data point acquired every 10 s, while CV was performed between − 0.8 V and + 0.8 V vs. the Pt RE using a scan rate of 10 mV s − 1. The electrochemical data were analyzed using ZView and CView software (Shribner Associates. Inc.). 3. Results and discussion 3.1. LSMC oxide composition The molar stoichiometries of the La1 − xSrxMn1 − yCryO3 ± δ (LSMC, x(Sr) = 0, 0.1, 0.2 and y(Cr) = 0.4, 0.5, 0.6) oxides, in their assynthesized form (calcined at 1260 °C for 12 h in air), were determined by inductively coupled plasma atomic emission spectroscopy (ICP-AES, Table 1), and were generally found to be quite close to what was expected from their co-precipitation synthesis. For the high Mn content LSMC oxides ((1−y)(Mn) ≥ 0.5), the Mn stoichiometries are a bit low, likely due to some uncertainty in the degree of hydration

of the Mn nitrate (Mn(NO3)2•xH2O, x ~ 4) precursor material, which in turn, affects the molar stoichiometry of La, Sr and Cr in these particular compounds (Table 1). The average Mn oxidation state of the La1 − xSrxMn0.6Cr0.4O3 ± δ (y(Cr) = 0.4) materials was determined by permanganate titration [14,15], thus giving their oxygen non-stoichiometry (3 ± δ, Table 1). This could only be achieved for the y(Cr) = 0.4 oxides, due to the limited solubility of the higher Cr content compounds in the Fe(II) solution (Fe(NH4)2(SO4)2•6H2O + 10% H2SO4). LMC64 (x(Sr) = 0, y(Cr) = 0.4) is seen to have an oxygen overstoichiometry of + δ = 0.17 in Table 1, which is higher than expected, based on the + δ = 0.08 reported for a similar compound, La0.974Mn0.584Cr0.390O3 [17], perhaps due to differences in the A-site deficiency. The oxygen over-stoichiometry in substituted LaMnO3 ± δ perovskites ((La, A’)(Mn, B’)O3 ± δ structure) arises from the Sr 2+ dopant (A’) at the A-site and the presence of Mn 4+ at the B-site, which is responsible for the H2 oxidation activity of these materials [18]. LSMC8264 (x(Sr) = 0.2, y(Cr) = 0.4) has an over-stoichiometry of +δ = 0.05, which is in very good agreement with literature values of + δ = 0.05 for a similar compound, La0.8Sr0.2Mn0.5Cr0.5O3 ± δ [18]. This previous work [18] also suggested that the oxygen overstoichiometry increases with increasing Sr content (x(Sr) = 0.2–0.3). This is observed for LSMC9164 (x(Sr) = 0.1), which has a lower over-stoichiometry of + δ = 0.03 than does LSMC8264.

3.2. Thermo-chemical stability of LSMC oxides in various environments Table 1 LSMC perovskite molar stoichiometry of La, Sr, Mn and Cr, and oxygen nonstoichiometry for the y(Cr) = 0.4 oxides. Name

3 ± δb

Stoichiometry Expected

Analysis by ICP spectroscopya La

LMC64 LSMC9164 LSMC8264 LMC55 LSMC9 155 LSMC8255 LMC46 LSMC9146 LSMC8246 a b

LaMn0.6Cr0.4O3 ± δ La0.9Sr0.lMn0.6Cr0.4O3 ± δ La0.8Sr0.2Mn0.6Cr0.4O3 ± δ LaMn0.5Cr0.5O3 ± δ La0.9Sr0.1Mn0.5Cr0.5O3 ± δ La0.8Sr0.2Mn0.5Cr0.5O3 ± δ LaMn0.4Cr0.6O3 ± δ La0.9Sr0.1Mn0.4Cr0.6O3 ± δ La0.8Sr0.2Mn0.4Cr0.6O3 ± δ

± 0.005–0.02. ±0.01.

0.92 0.86 0.72 0.96 0.86 0.81 0.99 0.90 0.79

Sr 0.10 0.20 0.097 0.81 0.11 0.21

Permanganate titration

Mn

Cr

0.66 0.65 0.67 0.52 0.51 0.51 0.42 0.41 0.43

0.43 0.41 0.43 0.52 0.54 0.51 0.60 0.59 0.59

3.17 3.03 3.05

3.2.1. (a) pXRD results The phase of the as-synthesized oxides (calcined for 12 h at 1260 °C in air) was determined in air by powder X-ray diffraction (pXRD), showing that most of these materials are rhombohedrallydistorted (R3c) perovskite-type oxides, as shown in pattern ‘a’ of Figs. 2–4 for LaMn0.6Cr0.4O3 ± δ (LMC64), La0.9Sr0.1Mn0.6Cr0.4O3 ± δ (LSMC9164), and La0.8Sr0.2Mn0.6Cr0.4O3 ± δ (LSMC8264), respectively. This air-stable structure has previously been reported in the literature for La0.75Sr0.25Mn1 − yCryO3 with varying Cr content [6,19]. Since the LSMC oxides are sintered on a yttria-stabilized zirconia (YSZ) disk at 1200 °C in air when fabricating half-cells for electrochemical testing, the stability of the LSMC powders was also investigated under these conditions. Comparison of these patterns with pattern ‘a’ (Figs. 2–4) shows that these materials retain their as-synthesized structure without any decomposition or phase change observed upon sintering in air.

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Fig. 2. Powder X-ray diffraction (pXRD) patterns of LaMn0.6Cr0.4O3 ± δ (LMC64) (a) as-synthesized (calcined 1260 °C for 12 h in air), and after (b) reduction at 800 °C for 120 h in H2 (3% H2O), (c) reduction at 1100 °C C for 24 h in 100% H2, and (d) re-sintering at 1260 °C in air after reduction at 1100 °C for 24 h. Note that peak labeled with (○) represents a La2O3 secondary phase.

The thermo-chemical stability of the LSMC materials was then evaluated under typical solid oxide fuel cell (SOFC) anode operating conditions, i.e., at 800 °C in humidified (3% H2O) H2 for 5 days (120 h) (Figs. 2–4, pattern ‘b’). The y(Cr) = 0.4 LSMC oxides undergo a transition to an orthorhombically-distorted perovskite structure upon reduction, which has previously also been observed for La0.75Sr0.25Mn0.5Cr0.5O3 in humidified H2 at 950 °C for 120 h [6,12]. This structural change is seen mainly in the primary peak at 31 2θ, where a doublet peak represents a rhombohedral distortion and a singlet peak an orthorhombic distortion [12]. For the y(Cr)= 0.5 and 0.6 LSMC oxides, with x(Sr) ≤ 0.1, an orthorhombic distortion is seen under both oxidizing and reducing conditions, similar to what has previously been reported for the x(Sr) = 0 LMC oxides [17]. For the x(Sr) = 0.2 LSMC oxides, all show a rhombohedral distortion under oxidizing conditions (e.g., y(Cr) = 0.4, Fig. 4, pattern ‘a’) and an orthorhombic distortion under reducing conditions (e.g., pattern ‘b’).

LSMC8264 (Fig. 4, pattern ‘b’) shows the formation of an additional phase, as seen by the new peak at 31.2 2θ (labeled as (•)), under these reducing conditions (800 °C for 120 h in H2 (3% H2O)). This is the primary peak of a Ruddlesden–Popper phase (R–P phase, a layered perovskite structure), similar to SrTiO3 [20], and consistent with the R–P phase ((La, Sr)2MnO4) formed when La0.75Sr0.25Mn0.5Cr0.5O3 ± δ is quenched from 900 °C in a reducing atmosphere (pO2 =10− 20 atm) [21]. A similar peak was also seen for La0.75Sr0.25Mn1 − yCr1 − yO3 ± δ (y(Cr) =0–0.6) after exposure to 5% H2:Ar at 950 °C for 24 h [19]. However, when y(Cr) ≥ 0.5, the LSMC oxides (x(Sr) =0.2) do not show the formation of this R–P secondary phase, confirming the reported stabilizing properties of Cr on LSMC oxides [19]. The LSMC oxides were exposed to 1100 °C in 100% H2 for 24 h (Figs. 2–4, pattern ‘c’) to determine the thermo-chemical stability of LSMC oxides under extreme reducing conditions. All LSMC oxides showed the formation of secondary phases, with the pXRD peaks

Fig. 3. pXRD patterns of La0.9Sr0.1Mn0.6Cr0.4O3 ± δ (LSMC9164) (a) as-synthesized (calcined 1260 °C for 12 h in air), and after (b) reduction at 800 °C for 120 h in H2 (3% H2O), (c) reduction at 1100 °C for 24 h in 100% H2, and (d) re-sintering at 1260 °C in air after reduction at 1100 °C for 24 h. Peaks labeled with (•) indicate the formation of a layered perovskite or Ruddlesden–Popper phase and (○) a La2O3 secondary phase.

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Fig. 4. pXRD patterns of La0.8Sr0.2Mn0.6Cr0.4O3 ± δ (LSMC8264) (a) as-synthesized (calcined 1260 °C for 12 h in air), and after (b) reduction at 800 °C for 120 h in H2 (3% H2O), (c) reduction at 1100 °C for 24 h in 100% H2, and (d) re-sintering at 1260 °C in air after reduction at 1100 °C for 24 h. Peaks labeled with (•) indicate formation of a layered perovskite or Ruddlesden–Popper phase.

labeled as (○) for La2O3 and (•) for the R–P phase. All LMC (x(Sr) = 0, y(Cr) = 0.4–0.6) oxides form La2O3 (e.g., Fig. 2, pattern ‘c’, y(Cr) = 0.4), characterized by a secondary phase primary peak at 29.7 2θ. The x(Sr) = 0.2 LSMC (y(Cr) = 0.4–0.6) oxides form a R–P phase secondary phase at 1100 °C in H2, with a primary peak at 31.2 2θ (e.g., Fig. 4, pattern ‘c’, y(Cr) = 0.4). The secondary phase formed by the x(Sr) = 0.1 LSMC oxides depends on the Cr content. When y(Cr) ≤ 0.5, both La2O3 and an R–P phases are formed (e.g., y(Cr) = 0.4, Fig. 3, pattern ‘c’), while when y(Cr) = 0.6, only the R–P phase is formed. This secondary R–P phase was seen only in Sr-doped LSMC oxides (x(Sr) ≥ 0.1), regardless of the Cr content, thus illustrating the destabilizing effect of a high Sr content, similar to what has been reported for La1 − xSrxMnO3 (LSM) materials [22]. The appearance of these secondary phases under extreme reducing conditions (24 h at 1100 °C in pure H2, Figs. 2–4, pattern ‘c’) suggests that the LSMC oxides may also form these phases over long periods of time at typical operating temperatures (700– 800 °C). Based on the pXRD results, we therefore expect these anodes to show gradual deactivating electrochemical behavior as a function of time under SOFC anode conditions, although it is not clear from the pXRD data which LSMC composition will be the least or most stable. The reversibility of the phase changes, seen after 24 h at 1100 °C in H2, was determined by re-sintering the reduced materials at 1260 °C for 12 h in air. Figs. 2–4 (pattern ‘d’) show that the reduced state secondary phases of LMC (Fig. 2, pattern ‘d’) and LSMC (x(Sr) N 0) oxides disappear and that the as-synthesized (Figs. 2–4, pattern ‘a’) perovskite structure is fully restored. These results (Figs. 2–4) confirm that the relative stability of the LSMC oxides under high temperature reducing conditions depends strongly on the Sr content, of relevance to longterm SOFC operation, even at lower temperatures (700–800 °C).

3.2.2. (b) Temperature programmed reduction (TPR) results To provide further guidance regarding which of the LSMC oxides would be the most stable as an anode, mass loss (TPR) analysis was carried out in 20% H2:80% between 150 and 1100 °C (1 °C min − 1). Fig. 5 shows the absence of any abrupt mass changes (stepped mass loss), suggesting that these materials are unlikely to undergo a sharp change in phase (and presumably, then, in H2 oxidation activity)

under typical SOFC anode conditions (i.e., low pO2, 600–1000 °C). Even so, two primary stages of mass loss can be discerned (Fig. 5), one at 350 to 450 °C (likely Mn 4+ to Mn 3+ reduction, Reaction (1)), and one at N450 °C (assumed to be Mn 4+ to Mn 3+ reduction, Reaction (2), and Mn 3+ to Mn 2+ reduction, Reaction (3)). A previous literature report for La1 − xSrxMnO3 [23] in a 10% H2:Ar atmosphere suggested that the reduction of Mn 4+ to Mn 3+ occurs at 150–530 °C [23] and the reduction of Mn 3+ to Mn 2+ at 550–930 °C, also reported previously [10] for the y(Cr) = 0.5 LSMC oxides. Another interpretation of the data in Fig. 5 is that the 350–450 °C mass loss is due to the reduction of only the Mn 4+ that is present in proportion to the oxygen over-stoichiometry (3 + δ), to Mn 3+ by Reaction (1), where the sub-script ‘segregated’ indicating that these species are isolated at the surface or grain boundaries of the oxide material. LSMC and LSM oxides [22,24] have similar oxygen nonstoichiometry (3 ± δ) behavior, such that when the oxide is over• stoichiometric, Mn 4+ (MnMn ) formation is compensated for by the formation of cation vacancies (V‴ La) (Reaction (1)). When the oxide is under-stoichiometric, Mn 4+ formation is compensated through the formation of anion vacancies (VO• •) (Reaction (2)). The N450 °C mass loss would then be related to one of two processes, depending on the amount of Sr present in the LSMC oxide. Thus, when x(Sr) = 0, the reduction of Mn 4+ to Mn 3+ by Reaction (2) could occur, while for the x(Sr) = 0.1 and 0.2 oxides, the reduction of Mn 3+ to Mn 2+ by Reaction (3) is proposed in this temperature range (N450 °C) [25]. ½LaðMn; CrÞO3 surface + VLa + VMn + 5Mn•Mn × × ×
LaLa + CrMn + 6MnMn bulk + 3H2 O •

x

x

x

x

•

+

3 H ↔ 2 2

ð1Þ

ð2Þ

••

ð3Þ

2MnMn + Oo + H2 ↔2MnMn + Vo + H2 O x

bulk

••

2MnMn + Oo + H2 ↔2MnMn + Vo + H2 O

2MnMn ↔MnMn + MnMn




ð4Þ

In Reactions (1)–(4), given in Kroger–Vink notation, represents × represents the Mn 3+ species, which is the Mn 4+ species, MnMn neutral with respect to the LSMC lattice, and Mn′Mn represents the

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Fig. 5. Temperature programmed reduction (TPR) of (a) LMC (x(Sr) = 0) oxides (y(Cr)= 0.4–0.6) and (b) La1 − xSrxMn0.5Cr0.5O3 y(Cr) = 0.5 oxides (x(Sr)= 0–0.2) obtained in 20% H2:80% He.

Mn 2+ species [10,14]. These reactions are derived from studies performed with the La1 − xSrxMnO3 perovskite oxide [22]. Either way, the reduction of Mn 4+ (Reactions (1) and (2)) would be expected to lead to a decrease in the electronic conductivity of the LSMC oxides, giving a significant loss of catalytic activity under reducing vs. oxidizing conditions [10]. However, it has been proposed for similar La1− xSrxMnO3 perovskites that they remain electronically conductive under reducing conditions due to the disproportionation of Mn 3+ (Reaction (4)), forming sufficient Mn 4+ species to maintain conductivity and anode activity. Indeed, we have previously shown that LSMC (y(Cr) = 0.5) oxide is electronically conducting, even under reducing conditions [7,10], supporting the suggested disproportionation process (Reaction (4)). Fig. 5(a) (x(Sr) = 0) shows that the LSMC materials with the highest Mn content (least Cr) are more reducible (greater% mass loss) in H2 than those with more Cr. This behavior was seen regardless of the Sr content (i.e., x(Sr) = 0–0.2). Similar results have been reported for La0.75Sr0.25Mn1 − yCryO3 as a function of Cr content as the temperature increases (350–1000 °C) in both N2 [6] and 5% H2:Ar atmospheres [18]. Reducibility, reported as % mass

loss [10,14], is correlated with the redox stability of these oxides, such that the least reducible oxides (lower% mass loss) are the most redox stable (Fig. 5). For the y(Cr) = 0.4 [14], 0.5 (Fig. 5(b)), and 0.6 (data not shown) materials, the Sr-free LMC oxides are the most reducible (Fig. 5(b)), followed by the x(Sr) = 0.2 and then the x(Sr) = 0.1 oxides, all evaluated at 800 °C. Ponce et al. (2000) have reported similar TPR trends, as a function of Sr content, for the Sr-doped (x(Sr) N 0) for La1 − xSrxMnO3 (LSM, x(Sr) = 0.1–0.5) in 10% H2:Ar between 25 and 1125 °C [23]. These trends are suggested to be related to the initial Mn 4+ content (oxygen non-stoichiometry) of each material, which was determined independently by permanganate titration (Table 1). Of all of the oxides studied here, LSMC9164 experienced the lowest mass loss between 150 and 1100 °C in 20% H2. Assynthesized LSMC9164 also had the lowest oxygen over-stoichiometry (3 ± δ) of the y(Cr) = 0.4 oxides (Table 1) and thus the lowest initial Mn 4+ content to be reduced. TPR reducibility trends (Fig. 5) have never been previously reported as a function of Sr content for the LSMC oxides, especially also including the x(Sr) = 0 LMC oxides (Fig. 5(a)). The correlation between the oxygen over-stoichiometry and reducibility data for the y(Cr) = 0.4 LSMC does suggest that these results are influenced by the initial Mn 4+ content at room temperature in air. However, either Mn 4+ or Mn 3+ reduction (Reactions (2) and (3), respectively) may occur during the N450 °C mass loss stage, depending on the Sr content of the oxide. If this is correct, then the % mass loss observed for x(Sr) =0.1 and x(Sr) = 0.2 LSMC oxides is indicative of the stability of the perovskite structure of these Srdoped LSMC oxides in a high temperature, reducing environment, which would not apply to the x(Sr) = 0 LMC oxides. The x(Sr) = 0.2 LSMC oxides show greater reducibility, compared to the x (Sr) = 0.1 LSMC, which may indicate that these oxides are forming more Mn 2+ (Reaction (3)), which would likely serve to destabilize the perovskite structure due to its larger size. For the similar La1 − xSrxMnO3 (x(Sr) = 0–0.5) oxides, the x(Sr) = 0.2 LSM perovskite structure has been reported to decompose (lose its perovskite structure) at a higher pO2 than seen for these x(Sr) = 0.1 LSM oxides at 800 °C [24]. The Sr-free x(Sr) = 0 LM and Sr-doped x(Sr) N 0 LSM oxides show very different behavior, and this is also observed for the LMC and LSMC oxides under study here. Here, the x(Sr) = 0 LMC oxides show more susceptibility to reduction, compared to the x(Sr) = 0.1–0.2 LSMC (e.g., y(Cr) = 0.5, Fig. 5(b)). This may reflect the potential for two different reduction processes at N 450 °C. These are the reduction of Mn 4+ , present in the oxygen understoichiometric (3 − δ) LMC oxide, to Mn 3+ (Reaction (2) in the case of LMC55), and the reduction of Mn 3+ to Mn 2+ (Reaction (3) in the case of LSMC9155 and LSMC8255). Overall, it is suggested that the TPR data should be interpreted separately for the Sr-free LMC (x(Sr) = 0) vs. the Sr-doped LSMC (x(Sr) = 0.1–0.2) perovskites.

Table 2 LSMC oxides (y(Cr) = 0.5) and their reducibility at 800 °C, as determined by TPR in 20% H2:80% He. Material a

LMC55 LSMC9155a LSMC8255a LMC(55)b LSMC(8255)b a b c d

A-sitec Non-stoichiometry

% mass lossd

0.96 0.96 0.99 099 1.0

6.8 3.8 5.0 13 1.6

This work, Fig. 5(a). Data from [10]. La and Sr molar stoichiometry determined by ICP-AES. Percentage mass loss at 800 °C in 20% H2:80% He, as determined by TPR.

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Based on the above discussion, it is clear that, for the y(Cr) = 0.5 perovskites (Table 2, Fig. 5(b)), the x(Sr) = 0.2 material is somewhat more stable in H2 (less% mass loss) than is the x(Sr) = 0 material. However, in our prior work [10] involving the use of a separately synthesized batch of LSMC materials, this trend was reversed. This may be related to the fact that the actual oxide compositions, of this previous batch [10], were more Mn-rich (La0.99Mn0.53Cr0.48O3 and La0.80Sr0.23Mn0.51Cr0.46O3, respectively, Table 2) and as the y(Cr) = 0.5 oxides studied here are generally more A-site deficient. These two factors may explain the differences in reducibility of LSMC (y(Cr) =0.5) oxides, in H2, observed in our two studies (Table 2) [10]. Consistent with this, it has been shown [26] that the extent of A-site deficiency in perovskite structures (e.g., La1 − xSrxTiO3) significantly affects their electronic conductivity, reducibility, and electrochemical performance. Other work has shown that the reducibility decreases as the degree of A-site deficiency decreases for Sr-doped (A-site) and Mn-doped (B-site) LaNiyFe1 − yO3 perovskite oxides in 5% H2:95% Ar between 150 and 800 °C [27]. Therefore, these relatively small, but significant, differences

75

in composition in the molar stoichiometries of La1 − xSrxMn0.5Cr0.5O3 oxides may account for the differences in our reducibility trends at 800 °C in 20% H2 (Table 2).

3.3. Electrochemistry of LSMC oxides in humidified H2 3.3.1. (a) General electrochemical performance To determine the H2 oxidation activity of the LSMC anodes under study here, the following sequence was used. First, the current was measured at an anode overpotential of 0.6 V for up to 4 hours, followed by cyclic voltammetry (CV) and open circuit electrochemical impedance spectroscopy (EIS) measurements, starting at 700 °C and then at 800 °C, all in humidified H2. Only the results for LSMC9164, determined by TPR to be the most stable anode, are shown in this section, as they provide a very good indication of the general behavior of all of the LSMC oxides. Notably, the catalytic activity of LSMC9164 (and the other oxides) was found to increase dramatically with increasing sintering time (10 h vs. 2 h) at 1200 °C. This is likely due to better adhesion of the anode layer to the YSZ

Fig. 6. Electrochemical response of LSMC9164, sintered at 1200 °C, at 700 °C in H2 (3% H2O). (a) Area corrected (0.4 cm2) i/t behavior at 0.6 V (vs. Pt reference), with each potentiostatic hold interspersed with CV and EIS experiments, (b) geometric area and area specific series resistance (ASRs) corrected CV response at 10 mV s− 1, and (c) EIS response at the OCP. Run numbers: (1) before polarization, and after holding at + 0.6 V for: (2) 2 min, (3) 4 min, (4) 8 min, (5) 16 min, (6) 32 min, (7) 64 min, (8) 128 min, and (9) 256 min.

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electrolyte, leading to better interconnectivity of reactions sites and thus sustained activity as a function of operating time. Fig. 6(a) shows that the current decreases with time at 0.6 V, then recovers significantly at the open circuit potential (OCP), but is lower in the next potentiostatic experiment. The gradual loss of performance of LSMC9164 may be associated with the loss of catalytic reaction sites or to changes in the solid-state properties of the LSMC oxide. A similar deteriorating response with time was seen at anodic (and cathodic) overpotentials for all of the oxides investigated in this work. Also, a higher anode activity (higher currents, lower resistance) is seen at 800 vs. 700 °C, as expected, for all of the LSMC oxides. The area specific series resistance (ASRs)-corrected CV response is shown in Fig. 6(b) for a fresh anode and after nine 0.6 V holding experiments. These results (for LSMC9164) are typical of what was seen for all of the LSMC oxides, at 700 and 800 °C, in this work. Notably, the cathodic branch (water reduction) shows characteristics typical of interfacial charge transfer reaction rate control, while the anodic branch (H2 oxidation) resembles the case of finite diffusion control [28]. It is interesting that these perovskite materials are somewhat more active towards water reduction than H2 oxidation. It is also seen that the CV currents (scans 1 and 9, Fig. 6(b)) are all quite similar, whereas the potentiostatic transients (Fig. 6(a)) show a much greater loss in activity over the same period. This argues that longer times at high overpotentials cause more degradation of anode performance. A similar CV shape to that seen in the anodic branch in Fig. 6(b) has previously been reported for commercial Ni-YSZ anodes in a directed flow cell geometry at 750 °C in H2 (3% H2O) between − 0.3 V and + 0.3 V, scanned at 20 mV s − 1 [28]. This CV response was ascribed to (finite) gas transport of H2O away from the triple phase boundary (TPB) of the Ni-YSZ working electrode (WE), using an isothermal, finite-gap stagnation point flow model [29] with reversible electro-

chemical reaction kinetics [28]. Under these conditions, a stagnant layer of fixed thickness is established above the WE under CV and constant current conditions, with a linear concentration gradient within this layer [29]. Based on Fig. 6(b), this interpretation may also apply to the H2 oxidation kinetics seen at LSMC oxides. Fig. 6(c) shows the Nyquist plots collected between the two CVs (Fig. 6(b)) at 700 °C in humidified H2. In all cases, the EIS data were collected from high to low (10 5 to 0.005 Hz) frequency and then in reverse. Since these spectra overlap, this demonstrates that no significant change has occurred in the electrode or in the H2 oxidation/H2O reduction reaction kinetics over the timescale of EIS data acquisition, indicating that the data can be reliably interpreted. It is also seen in Fig. 6(c) that, with time of polarization at 0.6 V, the overall area-specific polarization resistance (ASRp) and especially the resistance associated with the low frequency (LF) arc, increases, consistent with the results of Fig. 6(a). This may be due to a change in the Mn oxidation state under polarization, which does not fully recover with time at OCP, or to micro-structural changes in the WE. In contrast, the high frequency intercept (ASRs) does not change with time (Fig. 6(c)), arguing against the development of delamination or other interfacial problems during electrochemical testing. Fig. 7 shows the EIS data for two LSMC9164 replicate samples at 700 °C, as well as the overlay Nyquist plot generated from the best fit of the data to the model circuits shown in the inset to Fig. 7. Similar results (two or three time constants) were obtained for all of the LSMC oxides, both at 700 °C and at 800 °C (at which the ASRp values are all smaller). The presence of either two or three arcs is consistent with what has been reported [30,31] for the similar anode material, La0.75Sr0.25Mn0.5Cr0.5O3. Both circuits consist of a series resistance (R1 or ASRs) in series with two (Fig. 7(a)) or three (Fig. 7(b)) units of a resistor (R) and a constant phase element (CPE) in parallel (R-CPE). CPEs are represented by two variables, CPE-T and n, where n indicates

Fig. 7. OCP EIS response for two replicate half-cells (a and b) employing a LSMC9164 anode WE, sintered at 1200 °C and operating at 700 °C in humidified H2. Inset: Model circuits used to fit the EIS data.

L. Deleebeeck, V. Birss / Solid State Ionics 203 (2011) 69–79

the type of CPE [10]. An n ~ 1 is a pure capacitor (C), n = 0 is a resistor (R), and n ~ 0.5 is a diffusion control Warburg element (W). Here, subscript ‘4’ reflects the LF elements, subscript ‘3’ the intermediate frequency (IF) and ‘2’ the high frequency (HF) elements (Fig. 7(b)). ASRp is the summation of the R values of all of the impedance arcs. The impedance responses shown in Figs. 6(c) and 7 are generally quite similar to those seen for Ni-YSZ [32,33], but showing significantly higher ASRs and ASRp values [30,31]. This type of EIS response has been successfully modeled by a stagnant gas layer of finite thickness above the WE [32], consistent with the shape of the anodic CV branch in Fig. 6(b). This model was developed by Primdahl and Mogensen (1999) for a Ni-YSZ cermet anode operating in 97:3 H2:H2O at 1000 °C [32]. The HF arc was ascribed to the charge transfer reaction at the Ni-YSZ TPB, coupled with its double layer capacitance [32]. Bessler (2006) re-interpreted this model using a finite-gas stagnation point flow model [29]. The LF feature was attributed to bulk gas transport limitations at very low inlet gas flow velocity, involving a combination of diffusion and convection in the anode chamber. The IF feature was interpreted as being a Warburg-type higher frequency shoulder to the LF gas transport feature, and so is, itself, associated with the same gas transport process [29]. Jacobsen et al. (2008) have simplified this model using a continuously stirred tank reactor (CSTR) model, combined with the concept of diffusion in a porous electrode [33]. The IF feature is ascribed to Nernstian diffusion impedance from within the anode pores (diffusion in the pores), and the LF feature to the stagnant gas (gas concentration) layer above the porous anode under low flow rate conditions [33]. It is proposed that, for our LSMC oxide anodes, the interfacial H2 oxidation kinetics are correlated with the HF arc, diffusion in the oxide pores governs the IF arc, and the LF arc is due to gas concentration impedance [28,29,32,33].

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decreases as the quality of the contact between the Pt mesh current collector and the LSMC oxide WE improves [34], which should then also impact on the ASRp values. Indeed, in our previous work [10], the ASRs and ASRp values were similarly correlated and both were somewhat smaller than those in Table 3, due to the use of a higher mesh Pt gauze current collector (100 mesh [10] vs. 50 mesh here). In general, all of the LSMC oxides showed improved performance (lower ASRs and ASRp) at 800 °C (Table 3) vs. at 700 °C (e.g., LSMC9164, Figs. 6 and 7) in humidified H2, as expected. Table 3 shows that LSMC9164 exhibits the best catalytic activity (lowest ASRp at 800 °C) of the LSMC oxide compositions studied here. It should be kept in mind that all electrochemical tests reported in this work were performed using Pt mesh current collectors. Future work will include determining the catalytic activity of LSMC9164 when using other current collector materials (e.g., Au and Ag) to rule out any catalytic effects of Pt. The EIS response (Figs. 6(c) and 7) is characterized by a combination of resistors and CPEs (Table 3). In terms of the CPE values, the value of CPE(HF) (n(HF) ≤ 0.5), associated with the HF arc, is consistent with what is expected for charge transfer reactions in a porous electrode [33], while CPE(LF) and CPE(IF) are more typical of what is expected for transport controlled process(es). The low and intermediate frequency n values are both significantly less than 1, similar to what is expected of diffusion controlled processes in porous electrodes [29,35]. To more carefully examine the trends with LSMC oxide composition, Fig. 8 shows that, at a fixed Sr content, the ASRp values decrease as the Cr content decreases at both 700 and 800 °C. This suggests that LSMC materials with the highest Mn content have the highest catalytic activity towards H2 oxidation at SOFC anode operating temperatures, even though they are thermo-chemically less stable (Fig. 5(a)). In fact, it has been previously reported that it is the Mn content that imparts the catalytic activity to the anode, as demonstrated by van den Bossche et al. (2008) for CH4 oxidation at La0.75Sr0.25Mn1 − yCryO3 [6]. In terms of the catalytic activity of the LSMC oxides plotted as a function of Sr content, at fixed Cr content, Fig. 9 shows that, for both y (Cr) = 0.4 and 0.6, the x(Sr) = 0.1 oxides have the highest catalytic activity (lowest ASRp), followed by the Sr-free (x(Sr) = 0) LMC perovskites. For the y(Cr) = 0.5 materials (data not shown), the trends in ASRp depend on both the Sr content and operating temperature (Table 3). As seen in Fig. 9(a), LSMC8264 shows the poorest anode performance (highest ASRp) of the y(Cr) = 0.4 oxides. This lower catalytic activity may be related to the propensity of this oxide to form a non-reactive R–P phase with time at 800 °C in humidified H2 (Fig. 4, pattern ‘b’). Overall, LSMC8246 (y(Cr) = 0.6) has the lowest catalytic activity. However, it does not show the lowest overall thermo-chemical stability [14], though x(Sr) = 0.2 LSMC8246 is more reducible (higher% mass loss) than x(Sr) = 0.1 LSMC9146.

3.3.2. (b) Effect of LSMC oxide composition on H2 oxidation activity (from EIS data) In our previous work [10,14], we reported the ASRp values for the LSMC oxides after exposing them to a lengthy sequence of CVs, potentiostatic polarization experiments, and intermittent OCP impedance measurements. However, as shown in Fig. 6, the electrochemical activity does not easily reach a steady-state. Therefore, here, we report the average ASRp values for the LSMC oxides prior to any polarization studies. This should better represent the inherent activity of each perovskite material composition, and not the effect of a variable polarization history on material performance. The results were compiled from three replicate LSMC anodes of identical chemical composition that were electrochemically tested at the OCP prior to the polarization experiments, first at 700 °C and then at 800 °C. Table 3 shows the circuit parameters obtained, using the equivalent circuits shown in Fig. 7. It is seen that, for each oxide, the R(LF), R(IF), and R(HF) values are quite similar in value and that there is also a clear correlation between the ASRs and ASRp values. ASRs

Table 3 EIS-derived electrical parameters obtained for single phase LSMC perovskite anodes, sintered at 1200 °C for 2 to 10 h in air, at the OCP at 800 °C in humidified H2. Material

ASRsa

R(LF)a

CPE-T(LF) b(10− 2)

n(LF)

R(IF)a

LMC46 LSMC9146 LSMC8246 LMC55 LSMC9155 LSMC8255 LMC64 LSMC9164 LSMC8264

23 25 25 5.8 27 20. 9.2 4.2 17

58 61 270 6.1 22 28 5.7 2.1 30.

4.5 1.7 0.43 1.7 16 2.3 5.2 38 11

0.71 0.54 0.60 0.89 0.79 0.60 0.79 0.57 0.65

50. 24 130

a b

Ωcm2. Ssn cm− 2 [10,34].

CPE.(IF)b (10− 3) 2.8 2.5 1.7

n(IF)

R(HF)a

CPE.T(HF)b (10− 3)

n(HF)

ASRpa

0.48 0.62 0.38

78 19 83 24 18 32 7.9 1.9 22

0.0080 0.00018 0.0010 8.9 140 8.7 6.6 4.8 1.5

0.55 0.87 0.75 0.32 0.70 0.39 0.45 0.71 0.70

160 110 530 30. 58 60. 14 5.2 58

24

13

0.52

1.8 14

1.7 20.

0.77 0.59

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Fig. 8. ASRp (log scale) of La1 − xSrxMn1 − yCryO3 ± δ anodes, sintered at 1200 °C and operating at 700 and 800 °C in humidified H2, as a function of y(Cr) content: (a) x (Sr) = 0, (b) x(Sr) = 0.1, and (c) x(Sr) = 0.2.

By comparing the data in Fig. 9 with the TPR reducibility trends (Fig. 5(b)), it is seen that, for y(Cr) = 0.4, it is LSMC9164 that is the least reducible [14], and also the most catalytically active (lowest ASRp, Fig. 9). This is a very positive result, similar to what was shown in our previous work [10] for La0.99Mn0.53Cr0.48O3 vs.

La0.80Sr0.23Mn0.51Cr0.46O3 (Table 2). To ensure that these electrochemical differences are not due to differences in oxide microstructure or morphology, the LSMC oxide anode layers were examined by scanning electron microscopy (SEM) following halfcell testing. Fig. 10 shows the cross-sectional SEM images of the best performing LSMC9164 (a) and the worst performing LSMC8246 (b) electrodes following electrochemical testing, first at 700 and then at 800 °C, in humidified H2. This shows very similar particle sizes (~ 1 μm) and WE layer thicknesses (10– 20 μm) in both cases. This strengthens the argument that the differences in stability and activity of these two oxides are related to chemical composition differences (Sr and Cr content), and not to differences in morphology. Finally, all cross-sectional images revealed a similar degree of delamination upon half-cell unloading, regardless of electrode composition. These results demonstrate that the x(Sr) = 0.1 LSMC oxides show better redox stability, overall, compared to x(Sr) = 0.2 LSMC, resulting in a more stable perovskite structure under high temperature, reducing conditions and, hence, sustained catalytic activity towards H2 oxidation (low ASRp). For the Sr-doped LSMC oxides, we suggest that H2 oxidation is catalyzed by both the Mn 3+ and Mn 4+ sites, as previously proposed by Wan et al. (2006) [25]. However, the Sr-free (x(Sr) = 0 LMC) oxides show high catalytic activity (Table 3) due to the presence of a greater number of Mn 4+ active sites, compared to Sr-doped LSMC, which is corroborated by the oxygen over-stoichiometry of LMC64 (Table 1) and the TPR-measured reducibility of LMC oxides (Fig. 5(b)) [14]. Possible reaction mechanisms for the hydrogen and methane oxidation reactions on the Mn 4+ and Mn 3+ sites of LSMC-type anodes have been discussed in depth in the literature [6,23,25], and the possibility that surface Mn 3+ species serve as catalytic sites in LSMC anodes has been specifically noted [25]. Furthermore, it has been noted [6,23] that the presence of Mn, and especially Mn 4+ species, are related to the catalytic activity of these (and similar) perovskite catalysts. We therefore suggest the presence of two types of catalytic sites in the LSMC oxides, depending on the presence or absence of Sr in the lattice. One type arises from Mn 4+ and Mn 3+ sites present in Srdoped LSMC (e.g., LSMC9164), while the second involves Mn 4+ sites that are generated specifically by oxygen over-stoichiometry in Srfree LMC (e.g., LMC64). The presence of these two catalytic processes will be investigated in the future by determining the oxygen overstoichiometry, redox stability, and catalytic activity of x(Sr) = 0.05, 0.15 and 0.25 LSMC oxides (y(Cr) = 0.4). Also, future studies will focus on determining the effect of compositing the LSMC oxides with YSZ [15] and GDC, and the exposure of these composites to low ppm H2S in H2 environments [15,36], on the electrochemical activity. 4. Conclusions

Fig. 9. ASRp for La1 − xSrxMn1 − yCryO3 ± δ at 700 and 800 °C in humidified H2 at fixed Cr contents: (a) y(Cr) = 0.4, and (b) y(Cr) = 0.6, all sintered at 1200 °C, as a function of x (Sr) content.

La1 − xSrxMn1 − yCryO3 ± δ (x(Sr)=0, 0.1, 0.2 and y(Cr)=0.4, 0.5, 0.6) perovskite-type oxides were successfully synthesized by co-precipitation. Single-phase electrodes were studied to better distinguish the impact of the LSMC oxide chemical composition (Sr and Cr content) on H2 oxidation catalysis at 700 and 800 °C in humidified H2. All of the materials were found to be chemically stable in humidified H2 at 800 °C, except LSMC8264 (x(Sr)=0.2, y(Cr)=0.4), which formed a layered perovskite secondary phase under all reducing conditions. At higher temperatures (1100 °C in H2), the Sr-free LMC oxides were found to form an additional La2O3 phase, while Sr-doped LSMC materials formed a Ruddlesden– Popper (R–P) phase. The formation of these secondary phases under these conditions depends on the Sr content, but not on the Cr content, and was reversible for all materials upon re-sintering at 1260 °C in air. The oxygen non-stoichiometry (3 ± δ) of the low Cr content (y(Cr) = 0.4) LSMC materials was determined by permanganate titration, showing that Sr-doped LSMC (x(Sr)= 0.1) has the lowest + δ value. Low+ δ and a low initial Mn4+ content correlates with less mass

L. Deleebeeck, V. Birss / Solid State Ionics 203 (2011) 69–79

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Fig. 10. Resin-embedded cross-sectional scanning electron microscopy (SEM) images of (a) LSMC9164 (5000×) and (b) LSMC8246 (7000×) working electrodes (WE) following electrochemical testing.

loss during temperature programmed reduction (TPR) between 150 and 1100 °C in 20:80 H2:He. Here, the x(Sr) = 0.1 LSMC oxide materials were found to be the least reducible (lowest initial Mn4+ content) at high temperatures and thus likely the most stable SOFC anodes. Electrochemical experiments were carried out using a 3-electrode, half-cell configuration in humidified H2 at 700 and 800 °C, with all oxide anodes sintered at 1200 °C for 2 or 10 h. Cyclic voltammetry, potentiostatic holds, and electrochemical impedance spectroscopy (EIS) data, acquired at the open circuit potential (OCP), all show the presence of diffusion controlled processes. In the EIS response, the high frequency arc may be related to the interfacial H2 oxidation kinetics (charge transfer reaction), the intermediate and low frequency arcs to gas diffusion in the porous electrode and gas concentration changes outside the working electrode. Overall, a better electrochemical performance was seen at higher temperature for all LSMC oxides, as expected. The results show that the catalytic activity towards the H2 oxidation reaction, reported as the area-specific polarization resistance (ASRp), decreases as a function of increasing Cr content (y(Cr) = 0.4–0.6). Such an inverse relationship was not observed with respect to Sr content (x(Sr) = 0–0.2). The best performance (lowest ASRp) of the LSMC oxides studied here was demonstrated by LSMC9164 (y(Cr) = 0.4, x(Sr) = 0.1) and the poorest by LSMC8246 (x(Sr) =0.2, y(Cr)= 0.6) at both temperatures. Oxides with a higher Sr content (x(Sr) = 0.2) tended to show lower catalytic activity, vs. the x(Sr) = 0.1 oxides, perhaps due to their propensity to form secondary phases (Ruddlesden–Popper phases) during operation at 800 °C, as seen by powder X-ray diffraction (pXRD) studies of LSMC8264 (x(Sr) = 0.2, y(Cr) = 0.4). Sr-free (x(Sr) = 0) LMC oxides showed intermediate catalytic activity, compared to the x(Sr) = 0.1 and x(Sr)= 0.2 LSMC oxides. All of the oxide layers, on dense YSZ, had a very similar morphology and thickness, indicating that the differences in their performance arise from their compositional variations. Acknowledgements We gratefully acknowledge funding of this work by the Natural Sciences and Engineering Research Council of Canada (NSERC), as well as Conoco-Philips. This research was also supported through funding to the NSERC Solid Oxide Fuel Cell Canada Strategic Research Network from the NSERC and other sponsors listed online (www.sofccanada.com). We also acknowledge NSERC scholarship funding of LD and extend our thanks to Drs. J. Fournier, S. Paulson, H. Shahbaazi, R. Venkataramanayya, and S. Xia (University of Calgary) for technical assistance, as well as

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