Mo-Substituted Lanthanum Tungstate La - ACS Publications

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Mo-Substituted Lanthanum Tungstate La28−yW4+yO54+δ: A Competitive Mixed Electron−Proton Conductor for Gas Separation Membrane Applications M. Amsif,† A. Magrasó,*,‡ D. Marrero-López,§ J. C. Ruiz-Morales,† J. Canales-Vázquez,∥ and P. Núñez† †

Department of Inorganic Chemistry, University of La Laguna, 38200-La Laguna, Tenerife, Spain Department of Chemistry, University of Oslo, FERMiO/SMN, NO-0349 Oslo, Norway § Department of Applied Physics I, Laboratory of Materials and Surfaces, University of Málaga, 29071-Málaga, Spain ∥ Renewable Energy Research Institute, University of Castilla−La Mancha, 02071-Albacete, Spain ‡

S Supporting Information *

ABSTRACT: Molybdenum substituted lanthanum tungstate, La28−y(W1−xMox)4+yO54+δ (x = 0−1, y = 0.923), was investigated seeking for an enhancement of the n-type electronic conductivity for its use as a mixed electron−proton conductor in hydrogen gas separation membrane applications. The materials were synthesized by the freeze-drying precursor method, and they were single phase after firing between 1300 and 1500 °C for x ≤ 0.8. The crystal structure changed from cubic (x ≤ 0.4) to rhombohedral (x ≥ 0.6) with increasing the molybdenum content. Transmission electron microscopy (TEM) investigations revealed an ordering of the oxygen vacancies with increasing Mo-content, giving rise to superstructure domains. The dependency of the conductivity with the oxygen and water partial pressure showed that these materials are good mixed electron−proton conductors under wet reducing conditions for x ≤ 0.4. The conductivity of the materials with x ≥ 0.6 was dominated by electrons, and they are expected to be less chemically stable due to the lower redox stability of Mo6+. The total conductivities in humidified H2 were 0.016 S/cm for x = 0.2 and 0.043 S/cm for x = 0.4 at 900 °C, and they were stable under these conditions for more than 60 h. The ambipolar proton−electron conductivity was estimated to be ∼1.6 × 10−3 S/cm for x = 0.4 at temperatures as low as 600 °C, which makes this family of materials very interesting and competitive candidates for applications such as hydrogen gas separation membranes at lower temperatures than state-of-the-art materials. KEYWORDS: La28−y(W1−xMox)4+yO54+δ, La6WO12, La6MoO12, mixed electron−proton conductor, dense hydrogen permeable membrane WO3) is not a single phase.7,10,11 A solid solution was found for compositions with La/W ratio between 5.3 and 5.7 after firing at 1500 °C.7 Outside this compositional range, segregation of either La2O3 (La/W ≥ 5.8) or La6W2O15 (La/W ≤ 5.2) were found. These materials crystallize in a cubic structure with space group F4̅3m7 although recent investigations using synchrotron X-ray diffraction data have shown that the crystal structure deviates slightly from cubic.12 This latest study reports that lanthanum tungstate can be written more correctly as La28−yW4+yO54+δv2−δ, where tungsten partially occupies lanthanum sites, and that the amount of oxygen vacancies in the material is determined by the La/W ratio (v = 2 − δ; δ = 3y/ 2).12,13 The lanthanum ions are coordinated with eight oxygen

1. INTRODUCTION Rare-earth tungstates with a 3:1 Ln2O3−WO3 molar ratio have attracted attention due to their relatively high mixed proton− electron conductivity.1−8 The electrical properties of these materials were studied by Haugsrud et al.3 reporting that “La6WO12” (LWO60) is the rare-earth tungstate that exhibits the highest proton conductivity, dominating under a wet atmosphere below ∼750 °C, with a maximum value of ∼3 × 10−3 S/cm at 800 °C. At high temperatures, these materials are mixed ionic-electronic conductors exhibiting n- and p-type electronic conductivity under reducing and oxidizing atmospheres respectively, making them interesting as dense membranes for hydrogen separation. One of the main advantages of these materials is that they present high tolerance toward CO2, in contrast to alkaline-earth containing perovskites, based on thermodynamic data from Chentsov et al.9 Recent investigations showed that lanthanum tungstate with a La/W = 6 ratio (usually written as La6WO12 or 3:1 La2O3− © 2012 American Chemical Society

Received: June 4, 2012 Revised: August 15, 2012 Published: August 30, 2012 3868

dx.doi.org/10.1021/cm301723a | Chem. Mater. 2012, 24, 3868−3877

Chemistry of Materials

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supplied by Aldrich. La2O3 was precalcined at 1000 °C for 2 h to achieve dehydration and decarbonation. Stoichiometric amounts of La2O3 were dissolved in hot diluted nitric acid, whereas WO3 and MoO3 were dissolved in diluted ammonia. An ethylenediaminetetraacetic acid (EDTA) (99.5% Aldrich) solution was added as complexing agent in a molar ratio ligand:metal of 0.5:1. The different cation solutions were mixed, stirred for 15 min, and the pH was adjusted to 9 by ammonia addition, resulting in a homogeneous and transparent solution. The solutions were frozen by dropwise addition to liquid nitrogen, retaining the homogeneity of the original solution. The frozen drops were dehydrated by vacuum sublimation at a pressure of 1−10 Pa in a Heto Lyolab 3000 freeze-dryer during 2 days to produce the dried solid precursors. Such precursor powders were initially calcined at 300 °C to produce the organic matter pyrolysis. After that, the powders were ground and calcined again at 800 °C for 1 h to remove the residual organic species. These powders were pressed into pellets and fired between 1000 and 1600 °C for 15 min up to 3 h. Characterization of Powders and Sintered Specimens. X-ray diffraction patterns (XRD) were recorded using a Philips XPert Pro diffractometer, equipped with a Ge(111) primary monochromator with Cu Kα1 radiation (1.540 560 0 Å) and the XCelerator detector. The scans were collected in the 2θ range (5−100°) with a 0.016° step for 2 h. Structure refinements were performed using the FullProf and WinPlotr softwares.19 Electron probe microanalyses (EPMA) were carried out on the cross-section of dense bodies, polished down to 1/4 μm with diamond paste, in a Cameca SX100 spectrometer to verify the nominal composition and check for possible Mo and/or W evaporation. Selected area electron diffraction (SAED) and high resolution transmission electron microscopy (HRTEM) imaging were performed on a JEOL JEM 2011 electron microscope equipped with a (±20°) double-tilt sample holder, operating at 200 kV. Samples for TEM observation were prepared by dispersion of a very fine ground powder specimen onto a perforated carbon film supported on a Cu grid. HRTEM images were recorded using an Orius Gatan CCD camera (2 × 2MPi). Simultaneous thermogravimetric and differential thermal analysis curves (TG/DTA) were acquired with a Perkin-Elmer instrument (Pyris Diamond series) at a heating/cooling rate of 10 °C/min under wet (∼3% H2O) air. Dense ceramic pellets were obtained by uniaxial pressing the powders at 75 MPa and then sintered at 1400 °C for x ≤ 0.4 and 1500 °C for x ≥ 0.6 for 1 h. The morphology and grain size of the ceramics were observed using a scanning electron microscope (SEM) (Jeol LTD, JSM-6300). Electrical Characterization. Cylindrical pellets of 10 and ∼2 mm of diameter and thickness, respectively, were coated with Pt-paste on each side of the pellet and then fired at 800 °C for 15 min. Impedance spectra were obtained using a frequency response analyzer (Solartron 1260) in different dry and wet gases (Ar, O2, and 5% H2−Ar) in the 0.1 Hz to 1 MHz frequency range with an ac perturbation of 100 mV. The spectra were recorded upon cooling from 800 to 150 °C with a stabilization time of 30 min between consecutive measurements. Impedance spectra were analyzed with ZView program.20 The resistance and capacitance values of the different contributions were obtained by fitting the impedance spectra data with equivalent circuits. The equivalent circuit consisted of (RQ) elements in series, where R is a resistance and Q is a pseudocapacitance in parallel. The conductivity dependence with oxygen partial pressure and water vapor partial pressure was measured in a ProboStat measurement cell (NorECs, Norway) by the 2-point 4-wire method from 1000 to 500 °C under wet conditions (∼2.5% H2O). The conductivity was monitored versus time at each new set of conditions to ensure that equilibrium was achieved before taking a measurement. The conductivity in humidified H2 was monitored for samples with x ≤ 0.4 at 900 °C for 65 h in order to check the stability of the materials under reducing conditions.

atoms. The La(4a) [atom(Wyckoff position in the F43̅ m space group)] forms relatively regular LaO8 cubes, while La(24 g) has a more distorted environment. W(4b) sites are fully occupied and form octahedra with the surrounding oxygen atoms. An additional W(24 g) sits on the lanthanum site to form a stable solid state electrolyte,12 as shown in Figure 1.

Figure 1. Crystal structure of La28−yW4+yO54+δ (for y = 1).12

The apparent discrepancy between the first structural model (W sits in the middle of a cube, with partial oxygen occupancy)7 and the latter (W forms octahedra)12 has been related to the fact that the oxygen atoms around tungsten are disordered in space and time and average out as WO8 cubes. The cubic model is, therefore, a reasonable approximate description of an averaged high symmetry structure, although it does not describe the true local structure in detail. The aim of the present contribution is to obtain new compositions with enhanced electronic conductivity below ∼800 °C, while retaining similar proton conductivities, via the partial substitution of tungsten by a more reducible cation, such as molybdenum. Since the hydrogen flux in lanthanum tungstate at intermediate temperatures is limited by the electronic conductivity, this would increase the ambipolar conductivity and hydrogen permeation. A lower chemical and mechanical stability under reducing conditions may, however, be expected due to the higher reducibility of molybdenum compared to tungsten.14 Consequently, an optimization of the composition is required and will be investigated here. To the best of our knowledge, Mo-substitution in the La28−yW4+yO54+δ system has not yet been reported, but related compositions can be found in the literature, i.e. La2Mo2−xWxO9,15,16 Lu6Mo(W)O12.17 In the present paper, the synthesis of La28−y(W1−xMox)4+yO54+δ (y = 0.923) by the freeze-drying method is described. The crystal structure of the materials has been investigated by X-ray and electron diffraction. In addition, hydration characteristics and conductivity studies of the different compositions have been performed.

2. EXPERIMENTAL SECTION Materials Synthesis. For simplicity, the compositions La28−y(W1−xMox)4+yO54+δ (x = 0−1; y = 0.923) will be hereafter denoted as LWMx, where x is the molybdenum content. We have fixed the La/(W + Mo) ratio at 5.5 (y = 0.923) to avoid lanthanum oxide segregation based on an earlier report.7 Nanocrystalline powders of LWMx were prepared from a freeze-drying precursor method previously described.7,18 The starting materials used as reagents were the following: La2O3 (99.99%), WO3 (99.8%), and MoO3 (99.5%); all 3869



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Figure 2. XRD patterns of the series La28−y(W1−xMox)4+yO54+δ (0 ≤ x ≤ 1, y = 0.923).

3. RESULTS AND DISCUSSION Phase Formation. The XRD patterns after firing the precursors at different temperatures are shown as Supporting Information (Figures S1−S3). Samples with x ≤ 0.4 are crystalline at 800 °C and the XRD patterns show considerable peak broadening due to the nanometric nature of the powders (∼7 nm crystallite size). A minor amount of unreacted La2O3 is still present after firing at 1100 °C, which reacts completely above this temperature to form a single phase at 1300 °C. The materials with higher molybdenum content (x ≥ 0.6) require higher firing temperatures to obtain a single phase. Different phases are formed depending on the firing temperature for LWM0.6 and LWM0.8 compositions, and they are single phases after firing at 1500 °C. The composition without tungsten (LM, x = 1) can not be prepared as single phase even at temperatures as high as 1600 °C. Different lanthanum molybdates are found depending on the firing temperature of LM. Three phases are present at 1000 °C, which were identified as La2O3, La2MoO5, and La4MoO9. The rhombohedral phase, reported as La6MoO12, starts to form at 1400 °C and is the main phase at 1500−1600 °C. However, a lanthanum deficient impurity identified as La4MoO9 remained after firing at 1600 °C. No additional attempts were made to obtain single phase lanthanum molybdate. Figure 2 shows the XRD patterns of the series La28−y(W1−xMox)4+yO54+δ (y = 0.923) for 0 ≤ x ≤ 1. The compositions crystallize in the cubic system for x ≤ 0.4, while they are rhombohedral for x ≥ 0.6. The minimum sintering temperature also depends on the molybdenum content: 1300 °C for the cubic phases (lower Mo content) and 1500 °C for the rhombohedral phases (higher Mo content). It is likely that reducing the lanthanum content also decreases the temperature where the single phase can be formed for the latter compositions, in a similar fashion as reported for the Mo-free material earlier.7 Molybdenum and tungsten are relatively volatile elements, which could potentially lead to changes in the stoichiometry during the preparation. Therefore, microanalysis measurements have been performed to confirm the stoichiometry of the sintered specimens. From the results presented in Table 1, it

Table 1. Microanalysis Results Given by the EPMA Technique and Compared to the Nominal Ratios, Given in Atomic Percenta x

a

0.2

0.4

EPMA

nom

EPMA

nom

La/(W + Mo) La/W La/Mo x

5.63 ± 0.10 6.97 ± 0.16 0.24 ± 0.02 0.6

5.50 6.88 0.20

5.64 ± 0.10 9.37 ± 0.39 0.66 ± 0.08 0.8

5.50 9.17 0.67

EPMA

nom

EPMA

nom

La/(W + Mo) La/W La/Mo

5.60 ± 0.11 13.49 ± 0.61 1.41 ± 0.18

5.50 13.75 1.50

5.43 ± 0.07 27.2 ± 0.70 4.01 ± 0.14

5.50 27.5 4.00

The error in the measurements has been calculated as 2σst.dev.

seems that no significant tungsten or molybdenum losses occur in any of the investigated samples, in accordance with the results reported for the Mo-free composition.7 The measurements correspond to the result of the average of 10 different spots across the interface, to ensure reliability. Structural Analysis. The phases with the cubic symmetry (LWM0.2 and LWM0.4) were analyzed using the Rietveld method in the space group F4̅3m as reported by Magrasó et al. for LW.7 The distortions from the cubic structure for LW were only visible when using high resolution X-ray synchrotron radiation,12 and such small distortions are not detected using laboratory XRD data. Figure 3 shows the Rietveld refinement of LMW0.4. All the diffraction peaks can be indexed according to the F4̅3m space group, which indicates that there are no additional phases. During the refinement, the cation and oxygen content was kept constant. The cation and oxygen positions were refined, while the thermal factors for oxygen were also kept constant for a better convergence. The structural parameters for LWM0.2 and LWM0.4 are listed in Table S1 (Supporting Information). The agreement factors were RB = 6.11% and Rwp = 13.8% for LWM0.2 and RB = 7.06% and Rwp = 17.5% for LWM0.4. These are relatively high fit factors but still quite acceptable taking into account the rather large cell volume (∼1400 Å3). It should be noted that the thermal factors and 3870



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by XRD is relatively large due to the dominant scattering power of the heavy cations. Neutron diffraction data is therefore required for a more accurate determination of the oxygen parameters in these structures. The structural description reported here should be taken as an average structure which likely contains superstructure domains due to oxygen vacancy ordering as demonstrated below by HRTEM. It is worth mentioning that the cell volume increases with increasing Mo content in the cubic regime, drops with the phase transition, and then increases again in the rhombohedral regime (see Figure S4, Supporting Information). This is not expected, since the ionic radius of molybdenum (0.59 Å) is slightly smaller than tungsten (0.6 Å) in 6-fold coordination.21 The sample LM is not a single phase and does not follow this trend. It should be noted that a similar behavior has been reported earlier for La2Mo2−xWxO9 with increasing tungsten content.22,23 This anomalous volume change was related to the variation in the oxygen occupancy with the Mo/W content, leading to a lower coordination around tungsten/molybdenum, without a change in the cubic symmetry along the series. The refinement of the structure with higher Mo content was rather more complex due to symmetry loss. The complete crystal structure of La6MoO12 is not reported in the ICSD database, and hence, it can not be used as starting model for LWM0.6 and LWM0.8. The compound La6MoO12 was first synthesized in 1972,24 and it was considered to crystallize in a cubic structure similar to La6WO12. Later on, Cros et al.25 synthesized new phases in the system La2O3−MoO3 and the compound La6MoO12 was indexed as a rhombohedral cell (s.g. R3)̅ with the following parameters: a = b = 10.5417 Å and c = 9.8896 Å, which were taken as the initial parameters for our refinement. The main diffraction peaks could be indexed according to this cell, though the calculated Bragg positions of the smaller reflections are shifted compared to the experimental values. The XRD patterns were indexed again using the Dicvol program and the best results, with an agreement FORM factor

Figure 3. (a) Rietveld refinement obtained from XRD data for cubic LWM0.4 and (b) Le Bail refinement obtained from XRD data for LWM0.8, indexed as a rhombohedral unit cell. The small and nonindexed diffraction peaks are due to superstructure reflections.

occupation of oxygen sites were fixed for better convergence and that the error in the determination of the oxygen positions

Figure 4. (a−c) SAED patterns for the cubic LWM0.4 composition showing a view down the [111], [112], and [332] zone-axes. (d) HRTEM image revealing an incipient superstructure in the regions marked with arrows in the [111] zone axis. 3871



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Figure 5. (a−c) SAED patterns for rhombohedral LWM0.8 sample viewed down the [001], [11̅0], and [11̅1] zone-axes and (d) HRTEM image showing a 7 × 7 superstructure in the [001] zone axis.

cell with cells parameters of a = b = 28 Å and c = 9.81 Å. These results are in agreement with the XRD patterns as LMW0.8 can be described as a 7a′ × 7a′ × c superstructure (a′ = 3.9 Å). The unit cell volume of this compound is around 6600 Å3, and hence, a complete structural characterization of these new phases requires further studies by neutron and X-ray synchrotron diffraction. The HRTEM images confirm these results and the superstructure appears clearly in several zone axes (Figure 5d). From these results, one may assume that the gradual increase of the Mo content results in a larger amount of vacancies that become ordered giving rise to new structures. Moreover, the ordering of oxygen vacancies and-or the decrease in symmetry causes a marked decrease in the proton conductivity for the materials containing more molybdenum, as discussed further below. Thermogravimetric Analysis. The thermogravimetric measurements recorded under humidified air for all compositions showed a reproducible weight gain on cooling for both cycles. For clarity, only the curves taken on the last cooling for the different compositions are compared in Figure 6. It shows the typical behavior of a proton conducting material: the weight increases upon cooling, and this indicates the formation of protonic defects according to the hydration reaction. In acceptor-doped perovskites, the dopant is often charge compensated by the formation of effectively charged positive defects (e.g., oxygen vacancies, v•• O in Krö ger−Vink notation):

of 120, were obtained for a rhombohedral cell with the following cell parameters: a′ = b′ = 3.9845 Å, c = 9.8894 Å, and γ = 120°. It seems that a relation exists between these new cell parameters and those reported by Cros et al.25 The parameter c is almost identical in both cases, while a = √7a′. The new parameters allow the correct indexing of the most intense diffraction peaks. Unfortunately, they do not hold for the low intensity peaks; see the inset of Figure 3b. These low intensity reflections could in principle be related to the presence of small amounts of impurities but are more likely due to superstructure reflections, which are very common among rare-earth tungstates and molybdates. This will be assessed and confirmed by electron diffraction and HRTEM in the next section. Structural Analysis by Electron Diffraction. Selected area electron diffraction (SAED) performed on LWM0.4 reveals that the main reflections of the corresponding diffraction patterns can be indexed according to the cubic unit cell with a = 11.2 Å (s.g. F4̅3m) described above. Nevertheless one should note the presence of extra weak reflections and diffuse scattering along several directions (Figure 4), which may be ascribed to ordering of oxygen vacancies as usually described in fluorite-based systems with the formation of pyrochlore or C-type domains.26,27 This suggests that the distortion from cubic to rhombohedral structure detected by XRD for x = 0.6 occurs at a local level for lower Mo contents. Indeed, the corresponding high resolution images reveal the presence of nanodomains where an incipient superstructure can be detected (Figure 4d). Higher Mo content, i.e. for LWM0.8, causes a far more complex situation and the corresponding SAED patterns reveal that the weak diffuse reflections and scattering have evolved to sharp superstructure reflections as depicted in Figure 5. The SAEDs of the [110] and [111] zone axes show extra reflections at GF + 1/7 {110}*, which suggests the presence of a 7 × 7 superstructure in the ab plane of the cubic unit cell. Consequently, the SAEDs were indexed using a rhombohedral

x 2AO + Ln2O3 = 2AxA + 2Ln/B + v •• O + 5OO

(1)

and in such cases the hydration reaction of the oxygen vacancies and formation of protonic defects can be written as: H 2O(g) + OOx + vO•• = 2OHO•

(2)

Lanthanum tungstate is, however, a nominally undoped system, which contains a number of inherently deficient oxygen vacancies, and the hydration reaction can strictly not be written 3872



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Figure 6. Thermogravimetric curves performed under humidified air from 800 °C to room temperature. Only the last curves taken upon cooling are shown.

as in eq 2. The effective charge of all ions and the concentration of vacancies per unit formula depend on the exact La/W ratio according to the formula: La28−yW4+yO54+δv2−δ, (where δ = 3y/ 2). The hydration equation in such cases can be written using a compatible Kröger−Vink nomenclature for an inherently defective oxygen sublattice, described by Norby.28 This has recently been applied for lanthanum tungstate13 and the hydration reaction reads 4/56/ • 52/56 • H 2O(g) + O54/56O + v108/56 54/56O = 2OH54/56O

(3)

O(4/56)/ (54/56)O

where is the oxygen sitting on the oxygen site and acts (52/56)• as “the acceptor”, while v(108/56)• (54/56)O and OH(54/56)O are the oxygen vacancies and the proton defects in the structure, respectively. Given that the effective charges of the vacancies (∼1.93) and the protons (∼0.93) are very close to those expected from an acceptor-doped system (2 and 1, respectively), lanthanum tungstate behaves as if it was nominally acceptor doped. The reader is referred to a recent publication by Erdal et al.13 for more details. The weight uptake between 800 and 300 °C is ∼0.15% for LWM0.2 and LWM0.4, and somewhat lower for the compositions with higher molybdenum content, but still quite comparable to lanthanum tungstate.3,7,29 The dehydration temperature seems to decrease with Mo content, which would imply lower stability of the protonic defects with increasing Mo. Thermogravimetric measurements were performed in humidified 5% H2/Ar up to 800 °C, and they showed similar weight uptake compared to the oxidizing conditions reported above. This may imply that the materials do not reduce to a significant extent and are stable toward reduction up to 800 °C. In order to confirm the redox stability of the phases, all the samples with x ≤ 0.8 were thermally treated in 5% H2/Ar at 800 °C for 24 h. The XRD patterns after reduction did not show neither the presence of any additional phases nor amorphization of the structure, which indicates stability toward reduction up to that temperature. Sintering and Microstructure. The morphology of the sintered specimens is shown in Figure 7. No impurities or phase segregations at the grain boundaries were observed, in agreement with the XRD results. The composition x = 0.8, however, showed some small grains on the surface of the pellet, but these are likely due to a contamination during the sintering process.

Figure 7. SEM micrographs of the surface of (a) LW (LWO55) from ref 7; (b) LWM0.2 and (c) LWM0.4, sintered at 1400 °C for 2 h; and (d) LWM0.6, (e) LWM0.8, and (f) LM, sintered at 1500 °C for 1 h.

The relative densities of the specimens were higher than 95%. It is clearly observed from the micrographs that the grain size increases with molybdenum content and sintering temperature. The average grain size varies from 2.5 μm for LW to 3.5 μm for LMW0.4 sintered at 1400 °C, so the substitution of tungsten by molybdenum enhances the grain growth. A further increase of the average grain size is observed for samples with higher molybdenum content (and temperature), e.g. 13 μm for LWM0.6 sintered at 1500 °C. Nearly fully dense specimens can be obtained at 1400 °C, so freeze-drying is a suitable method to achieve dense bodies with the intended cation ratio. Electrical Characterization. The impedance spectra in the low temperature range show three separated contributions for LWM0.2 and LWM0.4 (see Supporting Information Figure S5). The equivalent circuit used at low temperatures can be expressed as (RbQb)(RgbQgb)Qe, where the subscripts b, gb, and e denote grain interior, grain boundary, and electrode, respectively. The typical capacitances of each process are in the order of picofarads per centimeter at high frequencies, nanofarads per centimeter at intermediate frequencies, and microfarads per centimeter at low frequencies. The grain boundary has a minor contribution to the total resistance of the electrolyte and is only visible at very low temperatures (below 350 °C) for LWM0.2 and LWM0.4. This is a major advantage when compared to other proton conducting materials, for instance Y-doped BaZrO3,30−33 which exhibits lower overall conductivities due to the large grain boundary resistance. The compositions with more molybdenum show only one arc in the 3873



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Figure 8. Dependence of the total (= grain interior) conductivity with the oxygen partial pressure for (a) LWM0.2 and (b) LWM0.4 between 500 and 1000 °C and (c) the comparison of the former with nonsubstituted LW (LWO54 from ref 13) at 800 °C to highlight the increase of n-type conductivity with Mo-substitution without a significant drop of ionic conductivity.

Figure 9. Dependence of the total (= grain interior) conductivity with water vapor partial pressure for LW (LWO57, i.e. La/W = 5.7), LWM0.2, and LWM0.4 between 500 and 800 °C under oxidizing conditions.

impedance spectra attributed to grain interior. Figure 8 shows the variation of the conductivity with the oxygen partial pressure (pO2) for LWM0.2 and LWM0.4 from O2 to H2 (2.5% H2O). The behavior is consistent with a mixed ionic-electronic conductor. The conductivity is independent of the pO2 at intermediate pressures, indicating dominating ionic conductivity. At high temperatures and high pO2, the conductivity increases with increasing the pO2, indicating the presence of ptype conductivity. At high temperatures and low pO2, the conductivity increases with decreasing pO2, which is consistent with n-type conductivity. The effect of Mo-substitution in lanthanum tungstate is clear: there is a strong enhancement of the n-type electronic conductivity for the compositions with higher Mo-content, while the p-type and ionic conductivity remains almost unaffected (see Figure 8c for clarity). This behavior is rather unusual and may be related to the special crystal structure of lanthanum tungstate compared to typical proton conductors with the perovskite structure. In perovskites, proton conductivity is typically enhanced via acceptor doping, while electron conductivity may be increased by donor doping, so that simultaneous acceptor and donor doping is difficult to achieve. On the other hand, proton conductivity in lanthanum tungstate occurs via an inherently deficient oxygen sublattice

and not through doping. The replacement of W by Mo may induce only a change in the electron conductivity, due to its higher reducibility compared to tungsten. The reaction of reduction of molybdenum in Kröger−Vink notation, and for a simplified “acceptor-doped” system, may be written as 2Mo xMo + OOx →

1 / O2 (g) + v •• O + 2MoMo 2

(4)

This implies that molybdenum reduces at low oxygen partial pressures, thus favoring the electron mobility through electron hopping between Mo atoms exhibiting different oxidation states. By increasing the Mo content, the number of charge carriers increases as well, and so does the electronic conductivity. For simplicity, we have chosen to write MoxMo (Mo(VI)) and Mo/Mo (Mo(V)), but other oxidation states for Mo may be more likely, e.g. Mo(IV). The dependence of the conductivity with water vapor in O2 is displayed in Figure 9 and compared to that of the Mo-free material with La/W ratio of 5.7 (LWO57), measured with the same equipment. The tested samples show a very similar trend. The conductivity decreases with decreasing pH2O, and the slope is steeper at the lowest temperatures as expected for a typical proton conductor. The flattening of the conductivity at 3874



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Figure 10. Arrhenius plots of the total conductivity for (a) LWM0.2, (b) LWM0.4, (c), LWM0.8, and (d) LM under various atmospheres.

low pH2O and high temperatures probably indicates that oxide ion conductivity becomes more dominating upon increasing the temperature due to the dehydration of the material, but a small contribution from electron holes may influence as well. LWO57 exhibits somewhat higher overall conductivities, likely due to the higher amount of expected oxygen vacancies, since the La/ (W + Mo) ratio is higher (5.7), compared to LWM0.2 and LWM0.4 (5.5). We may mention that it seems that there is a small, but measurable difference in the ionic conductivity between the compositions, in the order LWO54 > LWM0.2 > LWM0.4 (c.f. Figure 8c). This may indicate that molybdenum inhibits the hydration reaction (eq 2) slightly, which would be in good agreement with the general trend in water uptake detected by thermogravimetry. Regardless of this minor difference, the data set shows that the materials present similar responses, which may imply that they present a similar defect structure, comparable proton conductivities and σH+/σO2−. In order to confirm proton conductivity, D2O/H2O shifts were carried out in Ar (to minimize the influence of electrons and electron holes) at 600 °C. At this temperature, σH2O/σD2O was ∼1.3 for all samples (LWO54, LWM0.2, and LWM0.4), which confirms that proton conductivity dominate at this temperature. The variation of the conductivity pH2O under reducing conditions (Figure S6, Supporting Information) is smaller compared to oxidizing conditions and flatter for LWM0.4 compared to LWM0.2. This is expected from the responses versus the pO2 shown earlier (Figure 8): electron conductivity dominates under reducing conditions and the impact of pH2O on the conductivity is therefore smaller. The temperature dependence of the conductivity for the different compositions is displayed in Figure 10. The conductivity of all compositions increases significantly from dry to wet conditions, which supports the presence of proton conductivity. By comparing the conductivities of wet Ar (mainly reflecting ionsas extracted from the pO2 dependence in Figure 8) and wet reducing conditions (reflecting a contribution from both electrons and ions) for the whole series, we can conclude that the n-type electronic conductivity increases significantly with molybdenum content, as already mentioned above. The p-type electronic conductivity can be estimated by comparing the difference between Ar and oxidizing conditions (at equal pH2O). There is a small

contribution from electron holes to the overall conductivity at high oxygen partial pressures for all compositions, but with negligible differences with the increase of Mo content. This was already supported from the pO2 dependencies of LWM0.2 and LWM0.4 described above and does not seem to change with higher molybdenum content. The proton conductivity is very similar for LWM0.2 and LWM0.4 but decreases more than 1 order of magnitude at the lowest temperatures for LWM0.8 and LM (for clarity; see Supporting Information Figure S7). This may be rationalized by the decrease in crystal symmetry from cubic to rhombohedral as it is well-known that a decrease in symmetry can affect proton transport to a large extent33 and also the smaller cell volume for rhombohedral phases. In addition, the ordering of oxygen vacancies detected by SAED and HRTEM, which becomes more significant for high Mocontent compositions, may also explain the decrease in protonic/ionic conductivity. The lower conductivity of LM compared to LWM0.8 under reducing conditions (both indicating predominant electron conductivity) may be associated to impurities (LM was not single phase) and/or to possible degradation under 5% H2/Ar, due to the higher reducibility of Mo6+ compared to W6+. The compositions with higher Mo content are, however, less interesting from both conductivity and stability considerations, and were not further studied. The overall activation energy of the conductivity in wet Ar in the low temperature range (T ≤ 400 °C) increases with increasing Mo content (x): 0.61 (x = 0), 0.65 (x = 0.2), 0.71 (x = 0.4), and 0.77 eV (x = 0.8−1). It is expected that under these conditions, protons are the major charge carriers and that protonic conductivity dominates. These overall activation energies probably reflect the enthalpy of mobility of protons in these materials, in accordance with the one reported for the unsubstituted lanthanum tungstate (∼0.62 eV).13 The fact that the activation energy increases with Mo content may suggest that the mobility of protons is reduced with the introduction of Mo. One may speculate that the ordering of the oxygen vacancies is the responsible of this effect, but further investigations are needed to clarify this point. Under wet reducing conditions, the overall activation energy at similar temperatures decreases with increasing Mo content: 0.63 (x = 0−0.2), 0.50 (x = 0.4), 0.47 (x = 0.8), and 0.45 eV (x = 1). This 3875



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∼2 mL/cm2·min at 600 °C, using a gradient of pH2 of 1 order of magnitude, an ambipolar conductivity of ∼1.6 × 10−3 S/cm for, and assuming that bulk diffusion is limiting. One possible drawback of molybdenum substitution may be a lower chemical stability toward reduction, and therefore, it is important to evaluate whether the properties under operation conditions are stable over time. Consequently, the variation of the conductivity under reducing conditions (wet H2) is monitored as a function of time at 900 °C. Figure 11 shows

is in accordance with an increasing influence of electronic conductivity over the protonic one. In summary, the compositions with high Mo content (x ≥ 0.8) exhibit low protonic conductivities and are, therefore, dominated by electrons under reducing conditions between 200 and 800 °C, while the compositions with low Mo content (x ≤ 0.4) are mixed ionic−electronic conductors: protons at low temperatures and wet conditions and electrons (and oxide ions) at higher temperatures and reducing conditions. The compositions with x = 0.2−0.4 are very interesting for application as membrane materials for hydrogen gas separation, and the following section addresses the applicability of these materials. Applicability of the Studied Materials. Lanthanum tungstate and related compositions are new materials which have attracted considerable attention as dense membrane materials for hydrogen separation during the last years.3−5,34−39 At 1000 °C, the ambipolar proton−electron conductivity in lanthanum tungstate is approximately 10−3 S/cm,34 giving hydrogen fluxes comparable to the best mixed proton−electron conducting perovskites. The hydrogen flux of thick (∼1−2 mm) has been reported for several compositions, including LWO56 (La/W = 5.6) ∼0.04 mL/min cm2,34 “La5/6Nd1/6WO12” 0.046 mL/min cm2,36 and “Nd5/6La1/6WO12” 0.0145 mL/min cm2,37 all at 1000 °C. The hydrogen flux increases with increasing the temperature and is limited by n-type electronic conductivity below approximately 950−1000 °C.34 Above this temperature, protons are the limiting species due to the dehydration of the material. The replacement of tungsten by molybdenum in lanthanum tungstate is, therefore, very interesting and embraces a family of truly competitive materials from an application point of view, since the electronic conductivity was enhanced substantially without altering the protonic conductivity significantly. The ambipolar proton−electron conductivity can be estimated from the pO2 dependency reported above, i.e. ∼7 × 10−4 S/cm for x = 0.2 and ∼1.6 × 10−3 S/cm for x = 0.4, at temperatures as low as 600 °C. This is comparable, or even higher, than the ambipolar conductivity reported for LWO (∼10−3 S/cm at 1000 °C), though at a much lower temperature. This is a considerable improvement compared to the nonsubstituted LWO, which is a purer ionic conductor and the temperature must be relatively high to achieve sufficient electronic conductivity to reach sufficient ambipolar transport of both electrons and protons. Acceptor-doped SrCeO3 is the mixed electron−proton conducting oxide most studied so far.40−45 These materials are predominantly protonic conductors below 600−700 °C,40 so that high ambipolar conductivity can only be obtained via suitable doping and/or by increasing temperature. Oh et al.45 reported that the ambipolar conductivity of SrCe0.85Eu0.15O3‑δ was close to 10−3 S/cm at 900 °C, which is comparable to that of our materials at 600 °C. The hydrogen flux measured for 1 mm thick SrCe0.95Yb0.05O3−δ at 950 K is ∼0.007 mL/cm2min46 or ∼0.05 mL/cm2min for 0.8 mm thick SrCe0.95Tm0.05O3−δ at 950 °C.47 The flux can increase substantially with decreasing membrane thickness, as reported for a 33 μm thick SrCe0.7Zr0.2Eu0.1O3−δ (0.23 mL/cm2min at 900 °C),48 and for a 2 μm thick SrCe0.95Yb0.05O3−δ (∼15 mL/cm2min at 950 K).46 The main drawback for a widespread practical application of these materials is the low chemical stability, and the tungstates can offer a feasible alternative with increased chemical stability. The H2 permeation of a 5 μm-thick LMW0.4 is predicted to be

Figure 11. Evolution of the conductivity at 900 °C in wet H2 for LWM0.2 and LWM0.4.

that the conductivity remains constant for both LWM0.2 and LWM0.4 for ∼65 h at 900 °C, indicating stability of these compositions under these conditions. Additional stability tests were performed on powder samples annealed in 5% H2/Ar for 24 h at 800 °C, and analysis of the XRD patterns did not reveal any sign of degradation. All in all, the present report shows that both LWM0.2 and LWM0.4 exhibit mixed electron−proton conductivities, are stable materials under reducing conditions up to 900 °C, and represent very competitive candidates as materials for hydrogen gas separation membranes at temperatures as low as 600 °C, much lower than the typical mixed electron−proton conductors reported to date.

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CONCLUSIONS The replacement of tungsten by molybdenum in lanthanum tungstate La28−y(W1−xMox)4+yO54+δ is very interesting and embraces a family of truly competitive materials for applications as dense membrane materials for hydrogen separation. These materials show cubic structure for x ≤ 0.4 and rhombohedral superstructure for x ≥ 0.6. The electronic conductivity was enhanced substantially compared to the nonsubstituted material, without altering the protonic conductivity significantly. The ambipolar proton−electron conductivity is estimated to be ∼1.6 × 10−3 S/cm for x = 0.4 at temperatures as low as 600 °C, and the conductivities were stable over more than 60 h at 900 °C in wet reducing conditions.

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ASSOCIATED CONTENT

S Supporting Information *

Text and Figures SI-1−SI-7. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Telephone: +47-22840660. Fax: +47-22840651. 3876



Article

Notes

(30) Kjølseth, C.; Fjeld, H.; Prytz, Ø.; Dahl, P. I.; Estournès, C.; Haugsrud, R.; Norby, T. Solid State Ionics 2010, 181 (5−7), 268. (31) Babilo, P.; Uda, T.; Haile, S. M. J. Mater. Res. 2007, 22 (5), 1322−1330. (32) Duval, S. B. C.; Holtappels, P.; Vogt, U. F.; Pomjakushina, E.; Conder, K.; Stimming, U.; Graule, T. Solid State Ionics 2007, 178 (25− 26), 1437−1441. (33) Kreuer, K. D. Annu. Rev. Res. 2003, 33, 333−59. (34) Erdal, S. Hydrogen in Oxides: incorporation, transport and effects on electrical properties. PhD thesis, University of Oslo, 2011. (35) Weirich, M.; Gurauskis, J.; Gil, V.; Wiik, K.; Einarsrud, M.-A. Int. J. Hydrogen Energy 2012, 37 (9), 8056−8061. (36) Escolástico, S.; Solís, C.; Serra, J. M. Solid State Ionics 2012, 216, 31−35. (37) Escolástico, S.; Solís, C.; Serra, J. M. Int. J. Hydrogen Energy 2011, 36 (18), 11946−11954. (38) Czyperek, M.; Zapp, P.; Bouwmeester, H. J. M.; Modigell, M.; Ebert, K.; Voigt, I.; Meulenberg, W. A.; Singheiser, L.; Stöver, D. J. Membr. Sci. 2010, 359, 149−159. (39) Haugsrud, R. New high-temperature proton conductors for fuel cells and gas separation membranes. Handbook of Fuel Cells; Wiley International: New York, 2010; Vol. 5, pp 505−515. (40) Norby, T.; Haugsrud, R. Dense ceramic membranes for hydrogen separation. I: Nonporous Inorganic Membranes for Chemical Processing; Wiley-VCH Verlagsgesellschaft, 2006; p 1−48. (41) Song, S.-J.; Wachsman, E. D.; Rhodes, J.; Dorris, S. E.; Balachandran, U. Solid State Ionics 2003, 164, 107−116. (42) Qi, X.; Lin, Y. S. Solid State Ionics 2000, 130, 149−156. (43) Li, L.; Iglesia, E. Solid State Ionics 2003, 58, 1977. (44) Mather, G. C.; Saiful Islam, M. Chem. Mater. 2005, 17, 1736. (45) Oh, T.; Yoon, H.; Wachsman, E. D. Solid State Ionics 2009, 180, 1233−1239. (46) Hamakawa, S.; Li, L.; Li, A.; Iglesia, E. Solid State Ionics 2002, 148, 71−81. (47) Cheng, S; Gupta, V. K.; Lin, J. Y. S. Solid State Ionics 2005, 176, 2653−2662. (48) Li, J.; Yoon, H.; Wachsman, E. D. J. Membr. Sci. 2011, 381, 126−131.

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

We are grateful to Spanish MINECO for funding through Projects MAT2010-16007 and MAT2010-19837-C06-04 and INNOCAMPUS-UCLM and Junta de Andalucı ́a FQM-6680. Support from the nanoPCFC and HYDROX projects of the Research Council of Norway is also gratefully acknowledged.

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