Mutual Diffusion and Microstructure Evolution at the Electrolyte

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Mutual Diffusion and Microstructure Evolution at the Electrolyte
Anode Interface in Intermediate Temperature Solid Oxide Fuel Cell Zhi-Peng Li,*,† Toshiyuki Mori,† Graeme John Auchterlonie,‡ Yanan Guo,‡,§ Jin Zou,‡,§ John Drennan,‡ and Masaru Miyayama|| †

)

Global Research Center for Environment and Energy Based on Nanomaterials Science, National Institute for Materials Science, Tsukuba, Ibaraki 305-0044, Japan ‡ Centre for Microscopy and Microanalysis, The University of Queensland, St. Lucia, Brisbane, Queensland 4072, Australia § School of Engineering, The University of Queensland, St. Lucia, Brisbane, Queensland 4072, Australia Research Center for Advanced Science and Technology, The University of Tokyo, Tokyo 153-8904, Japan ABSTRACT: The microstructure and elemental distribution of the gadolinium-doped ceria (GDC) thin film electrolyte, Ni-GDC cermet anode, and the interface between them were comprehensively characterized by high-resolution transmission electron microscopy (HR-TEM) and energy-dispersive X-ray spectroscopy (EDX) operated in scanning TEM (STEM) mode. HR-TEM observations show newly appeared microstructure (i.e., superstructures) formations at both GDC and metallic Ni grains at the electrolyte
anode interface. STEM-EDX mapping and line scan analyses illustrate that not only can Ni diffuse into GDC grains as previously reported but also Ce and Gd can diffuse into metallic Ni particles with equal diffusion lengths as that of Ni diffusion. Such mutual diffusion is independent of ionic radii and can result in the valence state change of diffusing ions, verified by electron energy loss spectroscopy investigations. Therefore, the mutual diffusion and related microstructural evolutions are elucidated to be dominating factors that lead to the interfacial layer formation between anode and electrolyte, which is identified to have a considerable influence on the ionic conductivity behavior in intermediate temperature solid oxide fuel cells.

1. INTRODUCTION Solid oxide fuel cells (SOFCs), which can directly convert chemical energy into electricity and heat with high efficiency and little pollution as well as operate on various fuels, have attracted intense interest during the past decades.1
3 The typical operating temperatures for current SOFCs are around 800
1000 °C, which will lead to high energy loss, degrade cell components, cause phase instability, complicate sealing procedures, and maintain a high cost of cell materials.4,5 Therefore, lowering the operating temperature to the intermediate temperature (i.e., ∼500 °C) or even lower has been considered as an efficient way to solve such problems and as the trend for SOFC development. However, lower operating temperature will also result in increased voltage losses in SOFCs. To compensate for this loss, two possible ways are put forward: one is to use alternative electrolyte materials, and the other is to decrease the thickness of electrolyte to obtain lower ohmic resistance.5
8 The low ionic conductivity of conventional yttria-stabilized zirconia (YSZ), which has the general operating temperature above 800 °C,9 at reduced operating temperature becomes an obstacle to develop intermediate-temperature SOFCs (IT-SOFCs). The rare-earthdoped ceria are more favorable and considered to be a better candidate for ionic conductors that can be used in IT-SOFCs, due to their better thermal stability and higher oxygen ion r 2011 American Chemical Society

conductivity.10
12 In recent years, a large number of different rare-earth elements, such as Y,13 Sm,14 Gd,15 Ca,16 and Ho17 have been doped into ceria to improve the oxygen ionic conductivity, control microstructure, as well as modulate some physical properties. Among all the doped ceria, Gd-doped ceria (GDC) exhibits higher ionic conductivity and a better application potential.18 Another efficient way is to use thin film electrolytes. Various thin film deposition techniques have been developed for fuel cell applications, such as spray deposition,19 spin- and dip-coating,20 electrophoretic deposition (EPD),21 and so on. Among all these techniques, EPD has its advantages due to the low manufacturing cost, simple and fast deposition procedure, and more flexibility in the choice of substrate.22,23 In the EPD process, electrolyte materials are dissolved as colloidal particles. When applying an external electric field, the suspended particles will move to the substrate surface and coagulate as a dense layer. By carefully controlling the applied voltage and deposition time, thin films with various thicknesses can be deposited directly onto the substrate. Different electrolyte materials have been successfully Received: February 2, 2011 Revised: March 7, 2011 Published: March 23, 2011 6877

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The Journal of Physical Chemistry C deposited onto different substrates fabricated by the EPD technique, such as YSZ layers on a porous LaxSr1
xMnO3(δ (LSM) cathode,21 YSZ on NiO-YSZ substrate,24,25 and Smdoped ceria deposited (SDC) on NiO-SDC.26 Similar to alternative materials selecting for the electrolyte, for the cell sample preparation, another requirement needs to be satisfied for the accompanied anode. For example, the anode should be porous to allow gas permeation, while the catalyst is needed for fuel oxidation and should have both electronic and ionic conductivity, as well as obtain high density at the triple-phase boundary (TPB).27,28 Since metallic Ni can provide electronic conductivity and high catalytic activity, the mixture of Ni and doped ceria is widely used as the anode material, and the cermet anodesupported SOFC is also a potential technology for the IT-SOFC development.29,30 The premixed NiO will be reduced to metallic Ni in the anode atmosphere and result in a porous structure, not only for the fuel access to the TPB but also for the inhibition of coarsening of metallic particles. In this study, both optimized ways were used to prepare high-quality cell samples: GDC electrolyte thin film was fabricated by the EPD method on the NiO-GDC substrate, and an improvement of the cell performance may be expected. Co-sintering at high temperature (∼1400 °C) produces dense thin films and enhances the adhesion between the deposited film and the substrate. Even though the cell sample is for IT-SOFC application, hightemperature treatment is unavoidable in cell fabrication. Therefore, interdiffusion among cell components may occur, and subsequently, related changes in the composition, element spatial distribution, and microstructure are anticipated during fabrication processes.31 These changes should not be neglected when considering SOFC performance, especially when compositional and microstructural stabilities are some of the main concerns for SOFCs applications.32
34 Particularly, for thin film electrolytes, such changes will have a non-negligible influence for the electrical properties and conductivities.8 With a decrease in the electrolyte thickness, the interface becomes more important, and many reports have demonstrated that most of the macroscopic properties of cells were determined by the electrolyte
electrode interface.35,36 Thereby, developing high-quality electrolyte and anode interfaces, especially with long-term stability, is urgently required. Nevertheless, detailed understandings of interactions or changes in morphology or chemical composition among cell components are still insufficient. Even though most relevant reports have observed Ni diffusion across electrolyte or electrode layers,26,33 few studies have been carried out on the interdiffusion at the interfacial region between the electrolyte and electrode. For example, the diffusion and its relative pathway of Ni from the cermet anode to electrolyte, or other possible diffusion mechanisms occurring at the interface, have been rarely investigated. Our recent work demonstrated that the Ni diffusion, from the Ni/SDC or Ni/GDC cermet anode into the adhered SDC or GDC thin film electrolyte, can lead to interfacial layer formation at the interface.37,38 Moreover, it was found that Ni diffusion can enhance the microstructural inhomogeneity in the interfacial region, including nanosized domains and superstructures in GDC grains. This would consequently lead to a decrease in the ionic conductivity of the electrolyte film.37,38 However, due to the qualitative nature of the energy-dispersive X-ray spectroscopy (EDX) in the scanning electron microscopy (SEM),37,38 quantitative analyses about the interfacial layer as well as the detailed interdiffusion mechanism are necessary. Most of the previous observations of microstructures were detected at

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the electrolyte site and only focused on GDC grains; nevertheless, microstructure characterizations should not be neglected at anode site as well as Ni particles. Since previous studies solely focus on the Ni diffusion, the detection of other element diffusions and related microstructural evolution at the interface are still lacking. Furthermore, it is unclear whether such interdiffusion occurred before or after hydrogen gas treatment during cell fabrication. To address such issues, it is thereby essential to quantitatively characterize the interactions (e.g., diffusion) of all the constituent elements, composition distribution, and corresponding microstructural evolution and hence develop a detailed understanding of the interactions between component cell materials. This can help not only elucidate the interfacial layer formation at the interface but also obtain the fundamental understanding about diffusion mechanisms among SOFC cell components. Such an understanding can also shed light on the manufacturing process to achieve a high-quality electrolyte
electrode interface for IT-SOFC applications.

2. EXPERIMENTAL SECTION The 20 atom % GDC (20GDC) nanopowder was synthesized by the ammonium carbonate coprecipitation method.39 The synthesized precursor was calcined under flowing O2 gas (200 mL/min) at 800 °C for 2 h before it was used for the substrate preparation and thin film deposition. The anode substrate was prepared by mixing NiO and GDC nanopowders, which were ball-milled by using zirconia balls for 12 h. Mixed powders were compacted into cylindrical pellets by isostatically pressing under 200 MPa. The constituted nanopowders in suspension (mixture of acetylacetone, iodine, and 20GDC nanopowders) were deposited directly onto the NiO-GDC green body by the EPD method, and details about EPD description were reported elsewhere.40 By carefully modulating the deposition voltage and time (i.e., 10 V and 2 min in this study), the thickness and density of the thin film can be optimized. EPD-deposited pellets were cosintered at 1400 °C in air for 6 h, followed by heating at 600 °C for 2 h under flowing H2 gas (50 mL/min) to completely reduce NiO into metallic Ni. A NiO-GDC sample after thin film deposition was also prepared for comparison but without H2 gas treatment. For interfacial morphology observation, a cross-sectional TEM sample was prepared by sticking two identical EPD fabricated samples face to face and mounted on a Mo ring with a diameter of 3 mm. The assembled sample was mechanically ground, dimpled, and ion-milled in a precision ion polishing system to obtain electrontransparent thin areas. The morphology and microstructural features of nanopowders, sintered pellets, and EPD fabricated samples were initially observed by SEM (Hitachi S-5000). Detailed microstructures of thin film samples were characterized by high-resolution transmission electron microscopy (HR-TEM) and scanning TEM (STEM). HR-TEM observations and selected area electron diffraction (SAED) detections were performed on JEOL JEM2000EX, operated at 200 kV. To investigate the chemical composition, spatial distribution, and valence state change of constituent elements, EDX and electron energy loss spectroscopy (EELS) were conducted using a FEI Tecnai G2 F30 field emission gun (FEG) analytical electron microscope, equipped with a Gatan image filter and operated at 300 kV. The FEG with highly coherent and strong electron beam source can provide high-resolution EDX and EELS detections from nanometer-scale regions. In particular, the EDX investigation was performed in STEM mode with an extraction voltage of 4500 V. Calculations 6878

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Figure 1. SEM micrographs of (a) calcined nanopowder and cross-sectional morphologies of GDC thin film supported by (b) dense NiO-GDC and (c) porous Ni-GDC substrates, respectively.

Figure 2. BF TEM micrographs of electrolyte
anode interfaces of (a) unreduced (dense substrate) and (b) reduced (porous substrate) samples. (c) A typical nanosized domain (highlighted by a dashed line) observed in the GDC grain. The corresponding SAED pattern is shown in (d).

for quantitative STEM-EDX mapping and line scans were conducted offline using ES Vision v4.0.172. Since strong elastic scattered peaks may interfere with the EELS spectra, the target region was tilted off-zone to ensure that strong Bragg reflections can be avoided. Also, observation regions for EELS analysis were detected at ultrathin crystalline areas with a relative thickness of 0.2
0.3 mean free path to minimize multiple scattering. For high-resolution EELS, we lowered the extraction voltage to 3800 V so that less electrons are emitted, yielding a finer zero loss peak, and a better energy resolution can be obtained. The convergence angle and acceptance angle for EELS was about 0.5 and 5 mrad, respectively. Generally, each EELS spectrum was acquired within 2
5 s to minimize the electron beam induced radiation damage. The power-law method was used to subtract background noise for acquired EELS spectrum analysis, which was extrapolated by using the software DigitalMicrograph 3.9.0.

3. RESULTS 3.1. Microstructure Characterization. Highly dense sintered samples can be prepared from well-dispersed calcined nanopowders. Figure 1a is an example of well-dispersed 20GDC calcined nanopowders, with an average size about 30 nm. Such nanopowders were used for the fabrication of a dense electrolyte thin

film and anode substrate. The cross-sectional morphology of an IT-SOFC cell sample was initially observed by SEM. Figure 1b and c compare SEM results of interfacial morphology of unreduced and reduced samples. Note that dense GDC thin film can be well adhered onto the substrate by the EPD method. Due to the reduction of NiO to metallic Ni particles, the anode substrate becomes porous after H2 gas treatment, which can provide the suitable channels for fuel gas diffusing to the electrolyte. The detailed microstructures of thin film and substrate, especially the interface, were further characterized by TEM. Figure 2a and b are typical bright-field (BF) TEM micrographs of EPD fabricated samples before and after H2 reduction. The deposited thin film is homogeneous and dense, without any pinholes. The transformation from dense to porous substrate due to reduction is reflected in changes in the morphologies observed in TEM (Figure 2b), which is consistent with SEM results. HRTEM observations of GDC grains, at both thin film and substrate, show nanosized domains. Figure 2c is an example of a typical nanosized domain (marked by the dashed line) embedded in a ceria matrix at the anode, which has similar features as those widely observed in sintered bulk samples of rare-earth-doped ceria.41
48 For example, due to differences in composition and lattice, there is lattice distortion and reorientation between the nanodomain and neighboring ceria matrix. Moreover, significant diffuse scattering can be observed in the SAED pattern (Figure 2d, viewed along the ceria [110] zone axis (ZA)). These diffuse scattering features, other than those sharp diffraction spots coming from fluorite lattice, are attributed to inhomogeneous nanodomains in ceria and indicate the short-range ordered structures of nanodomains. Previous reports demonstrated that such nanodomain formation was due to the aggregation and segregation of dopant and oxygen vacancies and a higher oxygen vacancy ordering than the ceria matrix.44
46 Therefore, these nanodomains can act as traps or sinks and inhibit the diffusion of oxygen vacancies and lead to a decrease of oxygen ionic conductivity. Along with HR-TEM observations of GDC grains at the anode and electrolyte, microstructure investigations of both GDC and metallic Ni particles at the interfacial region are also important. It was found that there were some local structural modulations at the interface. Figure 3a is a HR-TEM image of one GDC grain observed at this area. A long-range ordering lattice feature can be seen, which is arising from the superstructure formation with a periodic lattice fringe distance of ∼0.81 nm. This distance is about five times that of the (113) lattice spacing of ceria (d{113} = 0.16 nm), verified by the corresponding optical diffractogram (inset in Figure 3a). Microstructure evolution occurred not only 6879

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Figure 3. (a) HR-TEM micrograph of one typical GDC grain with the superstructure formation. Inset is the corresponding diffractogram. SAED patterns were acquired from metallic Ni particles located at (b) the anode and (c) and (d) the electrolyte
anode interfacial region, respectively.

in GDC grains but also in metallic Ni grains. Figure 3b is the SAED pattern of a single grain observed at the interfacial region, which is quite similar to that of fluorite structure. By measuring the diffraction spot distances r1 and r2 as labeled in Figure 3b, which correspond to crystalline d-spacings of 0.204 and 0.125 nm separately, we can index the diffraction pattern and conclude that the phase is cubic Ni (a = 0.352 nm). Since it is pure metallic Ni, the nanodomains observed in GDC grains are not expected in such homogeneous grains. Due to larger particle size of metallic Ni than GDC grains, it is difficult to obtain electron-transparent thin areas in Ni particles exactly at the interface, and related HRTEM images are hence deficient in this study. However, microstructure modulations in metallic Ni can be revealed from corresponding SAED patterns. Figure 3c is a SAED pattern obtained from one single grain at the interface. With a similar diffraction pattern index as aforementioned, it can be determined as the metallic Ni and was viewed along the same [110] ZA as that in Figure 3b. Different from the SAED of pure metallic Ni in Figure 3b, extra periodic diffraction spots appeared (Figure 3c), which can be indexed as {hkl} þ 1/2{111} where {hkl} is the diffraction planes of cubic Ni. These extra diffraction spots are ascribed to microstructural modulations occurring in metallic Ni particles (i.e., the superstructure formation). Such local microstructure modulations can also be interpreted by SAED analysis along different ZA. Figure 3d is another SAED pattern of a metallic Ni particle observed along [112] ZA, with the extra diffraction spots also indexed as {hkl} þ 1/2{111}, indicating similar superstructure formation in the metallic Ni particle. This type of superstructure is different from those observed in 50 atom % Tb-doped ceria49 and 20GDC grains in the Ni-GDC sample,34,35 which were believed to be from inhomogeneous domains that existed in doped ceria and have a C-type-related structure. It thus implies that the constituent elements or even the formation mechanism of the superstructure observed at the interface region are uniquely different from previous reports.37,38,49

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Figure 4. STEM (a) BF and (b) DF images of the cross-sectional morphologies of reduced sample. (c) and (d) are the STEM EDX line scan results of the areas labeled by yellow and purple lines in (b). STEM EDX point scan results obtained at grains far away from the electrolyte
anode interface and at the interface are shown in (e) and (f), respectively.

To probe this problem, more detailed analyses, such as chemical compositions and related spatial distributions at the interface, are essentially needed. 3.2. Mutual Diffusion Between Metallic Ni and GDC Grains. The elemental distribution at the interfacial region was directly characterized by STEM-EDX. Figure 4a and b are crosssectional BF and dark-field (DF) STEM micrographs of a reduced cell sample, respectively. Besides a polycrystalline microstructure observed in the STEM BF image, there is the spatial distribution of different elements in the STEM DF image as well, indicated by the different contrast of different elements (e.g., Ni and GDC). From the random variation of contrast in the STEM DF image (Figure 4b), it implies that metallic Ni particles do not agglomerate and are homogeneously distributed in the anode. For comparison, investigations were initially performed at the region far away from the interface. Figure 4c is the STEM-EDX line scan across GDC grains at the region in the thin film electrolyte, as highlighted by the dashed yellow square in Figure 4b. There is no obvious elemental variation across the GDC grain interior and boundary. Ni was not detected at both grain interiors and boundaries at regions far from the electrolyte
anode interface. Similar STEM-EDX line scans were also conducted at the anode site (labeled by a solid green square in Figure 4b). Figure 4d illustrates that concentrations of Gd and Ce are uniform and show no fluctuation across grains or segregation at the grain boundary. According to the concentration profiles obtained from STEM-EDX detections, it suggests that there is no significant interdiffusion among all the grains at the region far from the interface. This was confirmed by the STEM-EDX point scan results. As shown in Figure 4e, all the peaks are assigned to constituent elements in detected grains. Only pure GDC and metallic Ni grains were detected. However, STEM-EDX point 6880

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Figure 5. STEM EDX elemental maps of Ni and GDC grains at the electrolyte
anode interfacial region. The examined elements are (a) Ni, (b) Ce, and (c) Gd. (d) Reconstructed map combining all three elemental maps.

scans of different grains at the interface region show different results. For example, in the metallic Ni particle, Ce and Gd peaks were observed (the black line in Figure 4f), while in GDC grains, Ni signals were detected (the red line in Figure 4f). This indicates that mutual diffusion may occur at the interface and result in the existence of Ni in GDC grains, and vice versa, Ce and Gd are anticipated in metallic Ni particles. Furthermore, it suggests not only that Ni can diffuse into GDC grain but also that Ce and Gd can diffuse into metallic Ni, which was rarely detected in previous reports.37,38 STEM-EDX elemental mapping was then employed to qualitatively investigate element spatial distribution at the interface. Figure 5a is the Ni elemental map acquired at the K-line spectrum of several contacted grains, which confirms that initial NiO has been reduced to metallic Ni. From Ni, Ce, and Gd elemental maps (Figure 5a to c), all the grain boundaries are not sharp; instead, diffuse boundaries exist among different adjacent grains. Especially, contrast gradients at grain boundaries in elemental maps reflect the changes of element distribution, for example, the decreasing concentration from one grain to its neighbor. Recombining the Ni, Ce, and Gd elemental maps, the trace of element distribution at grain boundaries is demonstrated in Figure 5d. This represents interdiffusion zones which exist at grain boundaries, with mixed elements diffusing into neighboring grains. To quantitatively evaluate the interdiffusion occurring among different grains at the electrolyte
anode interface, STEM-EDX line scans were conducted. Since the focus is the interdiffusion between the Ni particle and GDC grain, first, STEM-EDX point analysis was used to determine the composition of different grains (data are not shown here). Subsequently, a STEM-EDX line scan was acquired vertically across the grain boundary of two nearby grains. For a comprehensive study, similar STEM-EDX line scans were also performed in unreduced cell samples. Figure 6 compares typical STEM-EDX line scan results conducted at the interface of both reduced and unreduced samples. The line scan results clearly illustrate the elemental distributions as a function of distance from the interface. The

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Figure 6. Concentration profiles of STEM EDX line scans conducted at the electrolyte
anode interface of reduced samples ((a) and (b)) and unreduced samples ((c) and (d)), respectively.

interface in STEM-EDX results is assigned as the intersection of concentration profiles of all constituent elements. Figure 6a and b are STEM-EDX line scan results obtained from the reduced sample. Note that the concentration of Ni varies considerably with the distance from the edge of the Ni particle. Similarly, the concentration variation can be detected at the nearby GDC grain. The concentrations of Ce and Gd decrease gradually across the interface and fade away in the adjacent Ni particle. Therefore, all the Ni, Ce, and Gd commix and mutually diffuse into neighboring grains, and the mutual diffusion zone is highlighted in blue. Similar mutual diffusion between Ni and GDC can be widely detected at the interface of reduced samples, and one more typical result is shown in Figure 6b. Statistic STEM-EDX line scan results demonstrate that the average mutual diffusion length is around 200 nm. According to these investigations, it reveals that all the constituent elements can mutually diffuse into adjacent contacted grains. Moreover, diffusion lengths of Ni, Ce, and Gd are almost equal and independent of ionic radii. Additionally, this type of mutual diffusion is not a global phenomenon but highly localized on the nanometer scale at the electrolyte
anode interfacial region. Comparable STEM-EDX line scans were also performed at the interface of unreduced samples, and corresponding concentration profiles are shown in Figure 6c and d. From Figure 6c and d, it can be noticed that concentrations of Ni, Ce, and Gd decrease dramatically at the interfacial region. Comparing STEM-EDX results in reduced and unreduced samples, it reveals that mutual diffusion zones in unreduced samples are much narrower than those in reduced samples. Quantitative results display that the average mutual diffusion length in unreduced samples is less than 50 nm. Considering the actual width of the grain boundary, due to the random orientation of polycrystalline grains, the diffusion length detected in unreduced sample can be neglected. According to the investigation shown in Figure 6, it can be concluded that mutual diffusion mainly occurs at the interface of reduced samples. It also indicates that metallic Ni can diffuse easier or faster than ionic Ni in the oxide. Particularly, Ni, Ce, and Gd have similar probability to mutually diffuse into nearby contacted grains reflected by similar detected diffusion lengths. 6881

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Figure 7. EELS spectra of Ce M4,5 adsorption edges acquired at the (a) electrolyte, (b) electrolyte
anode interface, and (c) anode separately. (d) EELS spectrum of the Ni L edge acquired in the metallic Ni particle at anode.

3.3. Valence State Change Detected by EELS. Due to the mutual diffusion between metallic Ni and GDC grains at the interface, the diffusion zone will form in terms of the mixture of diffused ions. It is anticipated that the valence state of constituent elements may change at the mixed ionic environment. To verify this suspicion, it needs to examine the detailed valence state or the related evolution at the mutual diffusion zone. EELS spectra were then acquired in the reduced sample since the mutual diffusion zone is negligible in the unreduced sample based on the aforementioned STEM-EDX line scan analyses. The valence state of each element can be determined from fine structures in EELS spectra. For example, the M4 and M5 adsorption edges of Ce can provide necessary information regarding related oxidation states.50 Figure 7 a
c display the EELS spectra of fine structures of Ce M4,5 adsorption edges acquired at the electrolyte, interface, and anode separately. Two featured twin peaks, Ce M4 and M5, can be clearly distinguished. The EELS spectra, acquired at the electrolyte and anode, illustrate that the predominant Ce is in a tetravalent state. This is determined by the similar satellite structures of postedge peaks that appear in the M4,5 peaks (denoted by two arrows in Figure 7a). The relative intensity ratio of M5/M4 can be used for the quantification of Ce3þ and Ce4þ. Such a type of fine structure has been well established as a fingerprint for the valence determination in EELS spectra analysis. Interestingly, the fine structure of the Ce M4,5 adsorption edge is quite different at the interface; for example, the postedge peaks of M4,5 disappeared. The relative intensity ratio of M5/M4 significantly changed, from the original ∼0.90 (Figure 7a and c, stronger Ce M4 peak) to ∼1.16 (Figure 7b, stronger Ce M5 peak). This fine structure change clearly indicates that the trivalent Ce is the predominant state at the diffusion zone. In addition, comparing the EELS spectrum of the metallic Ni L edge (Figure 7d), weak Ni L edge peaks can also be observed in the EELS spectrum acquired at the mutual diffusion zone (denoted by the arrows in Figure 7b). Even though the exact Ni valence state can not be determined from such weak

peaks, it implies that Ni has a valence state change due to the slight change of relative positions between Ni L2 and L3 peaks. This observation is in accordance with STEM-EDX results that it is a mixture of diffused ions at the mutual diffusion zone. As a result of mutual diffusion among Ni, Ce, and Gd, it can be expected that the concentration of oxygen vacancies should increase and may result in an enhanced oxygen vacancy ordering. This hypothesis can be verified by the EELS investigation of the fine structure of oxygen K edge peaks, which has been used to detect the oxygen vacancy ordering level in rare-earth-doped ceria.44,45 Figure 8a
c are EELS spectra of the oxygen K adsorption edge, acquired at the electrolyte, interface, and anode, respectively. The fine structures, such as typical featured peaks A, B, and C, can be clearly distinguished from the spectra acquired at the electrolyte and anode (Figure 8a and c). However, the fine structure changed in the spectra acquired at the interface, with the combination between peaks A and B. Moreover, the relative intensity ratio of peak B and C (IB/IC) varies from ∼0.75 (Figure 8a and c) to ∼0.94 (Figure 8b). This type of fine structure evolution was applied as the fingerprint of determining the oxygen vacancy ordering level in nonstoichiometric oxides.51 Therefore, based on EELS spectra comparison, an enhancement of oxygen vacancy ordering is determined in the mutual diffusion zone. Similarly, EELS spectra were also acquired for the Gd M4,5 adsorption edges at the electrolyte, interface, and anode separately (Figure 8d
f). From the comparison of fine structures of Gd M4,5 peaks, there is no obvious variation among all the Gd spectra acquired. It hence indicates that no valence state change occurred for Gd diffusion ions, which is different from that of diffused Ni and Ce ions.

4. DISCUSSION STEM-EDX and EELS studies allow us to evaluate the diffusion processes and related evolution of microstructural features observed at the electrolyte
anode interface. On the 6882

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Figure 8. EELS spectra of oxygen K edges acquired at the (a) electrolyte, (b) electrolyte
anode interface, and (c) anode separately. Gd M4,5 edge spectra acquired at the electrolyte, electrolyte
anode interface, and anode are shown in (d), (e), and (f), respectively.

basis of these results, it is assumed that there is a certain intrinsic relation among the different microstructural features, element spatial distribution, as well as valence state evolution. The STEM-EDX investigations clearly demonstrate that the mutual diffusion, between metallic Ni particles and GDC grains, mainly occurred at the electrolyte
anode interface in reduced samples. Nevertheless, such mutual diffusion was rare in unreduced samples. Since metallic Ni particles are usually considered as large surface area and high surface energy solid, there is an intrinsic thermodynamic driving force to minimize the total free energy, for example, by reducing surface area or enhancing the probability of diffusion.52 It is hence reasonable to expect that metallic Ni can diffuse easier and faster than that of Ni2þ in NiO. Additionally, the calculated phase stability diagram and experimental results have established that NiO does not react with CeO2, and there is no solubility of NiO in ceria.53
55 Therefore, the solid solubility of NiO in CeO2 can be negligible. It is thus believed that the mutual diffusion is mainly driven by the diffusion of metallic Ni, rather than high-temperature sintering, since both samples had been sintered at the same high temperature (1400 °C). Moreover, it was verified that not only Ni can diffuse into GDC grain, which is in accordance with previous reports,37,38 but also Ce and Gd can diffuse into metallic Ni particles. Additionally, both Ni and Ce and Gd have equal diffusion lengths. Considering the large difference in the radii between Ni (0.124 nm) and Ce4þ (0.087 nm), or Ni and Gd3þ (0.094 nm), this type of mutual diffusion should be independent of radius and can be considered as substitution diffusion. Furthermore, during H2 gas treatment for the reduction of NiO to metallic Ni, Ce4þ will be partially reduced to Ce3þ simultaneously. Due to a smaller radius difference between Ce3þ (0.102 nm) and Ni, Ce3þ can be considered as a suitable substitute for the Ni diffusion and dissolution in GDC grains because of the decreased valence state difference. This can also enhance the mutual diffusion and lead to the mutual diffusion zone mainly formed in the reduced sample, instead of unreduced NiO-GDC samples. Due to the mutual diffusion between metallic Ni and GDC grains, the formation of a diffusion zone will consequently appear at the interfacial region. This kind of diffusion zone was regarded

as a mixture of diffused ions. EELS analyses verified that the diffused Ni and Ce had valence state change (e.g., the metallic Ni transits to the Ni2þ cation). The dissolution of Ni in ceria can be described through the following empirical reaction CeO2

00

NiO sf Ni

Ce

þ V O :: þ OO 

where the Kr€oger
Vink notation is used. This will result in the formation of abundant oxygen vacancies at the diffusion zone. Meanwhile, the substitution diffusion is expected to be an easier way to accelerate the mutual diffusion since diffusing elements have significant differences in ionic radii. For the balance of local oxygen vacancy concentration, Ce4þ cations need to be reduced to Ce3þ and substituted diffuse into the metallic Ni particle. The diffusion of Ce3þ cations will also associate with the diffusion of nearby dopant Gd3þ cations. Therefore, mutual diffusion can enhance the instability of cell components and affect oxidation states of cell composition. For example, the reduction of Ce4þ to Ce3þ at the interface region will result in a mixed ionic/electronic conductivity and a decreased open circuit voltage.28 Therefore, the mutual diffusion and related formation of diffusion zone are the dominant factors that lead to the microstructural feature changes (i.e., superstructure formation) at the electrolyte
anode interfacial region. On the other hand, the newly appeared superstructure can provide some possible ways for the assumed trapping effect or sinks of oxygen vacancies56,57 and correspondingly retards the transport of oxygen vacancies. Consequently, it will lead to the enhancement of oxygen vacancy ordering at the mutual diffusion zone, identified by the evolution of oxygen K edge peaks detected by EELS. It also implies that the oxygen diffusion across the interface may be significantly hindered due to the formation of a superstructure, which may strongly increase the relative resistivity.58,59 Due to the enhanced inhomogeneity, the formation of an interfacial layer is considered as a high resistance layer at the electrolyte
anode interface, which will directly reduce the conductivity and degrade the SOFC performance. This also provides a possible explanation for the reduced conductivity 6883

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The Journal of Physical Chemistry C measured in the thin film electrolyte cell samples.37 Note that such type of microstructural change and negative interfacial layer can even occur before operating processes, which draws attention to improve high-quality cell preparations.

5. CONCLUSIONS GDC electrolyte thin film supported by the Ni-GDC cermet anode was fabricated by an EPD technique. Microstructural features of cell samples, especially at the electrolyte
anode interfacial region, have been studied by SEM and HR-TEM. Short-range ordering inhomogeneities (i.e., nanosized domains) were observed in GDC grains at both electrolyte and anode sites. However, microstructure evolution occurred at the interfacial region, and long-range ordering inhomogeneities (i.e., superstructure) began to appear. On the basis of STEM-EDX investigations, it identifies that mutual diffusion mainly occurs at the interface. Different from previous results, not only Ni can diffuse into GDC grains but also rare-earth elements Ce4þ and Gd3þ can diffuse into metallic Ni particles. Moreover, such kind of mutual diffusion is independent of ionic radii, and all the diffused ions have similar diffusion lengths. This leads to the formation of a diffusion zone and provides the possible reason for the superstructure formation at the interface. Additionally, it was verified that such mutual diffusion only occurs in the reduced sample, indicating that only metallic Ni can diffuse instead of Ni2þ in NiO. EELS analyses demonstrate that the diffusion zone can be considered as a mixture of diffused ions with changed valence states. The valence state change of Ni and Ce is believed to be a dominant factor that results in the mutual diffusion and superstructure formation. This can consequently enhance the oxygen vacancy ordering level and may provide the reasonable explanation for the decreased conductivity as previously reported.17 Therefore, the evolution of composition and microstructures should not be neglected even before cell operation, to develop high-quality cell components for IT-SOFCs application. ’ AUTHOR INFORMATION Corresponding Author

*Tel.: þ81-29-8513354, ext. 8544. Fax: þ81-29-8604712. E-mail: [email protected].

’ ACKNOWLEDGMENT The financial support from the Grant-in-Aid for Scientific Research (22310053) by the Ministry of Education, Culture, Sports, and Technology (MEXT), Japan, is gratefully acknowledged. The authors also appreciate the partial funding support from Global Research Center for Environment and Energy based on Nanomaterials Science (GREEN), National Institute for Materials Science, Japan. Z.-P. Li thanks Prof. Fei YE (Dalian University of Technology, China) and Dr. Dingrong OU (Dalian Institute of Chemical Physics, China) for helpful discussions. Moreover, G. J. Auchterlonie is grateful for financial support from the Australian Academy of Sciences - Japan Society for the Promotion of Science Bilateral Exchange Program, Scientific Visits to Japan (2010-2011-RC 21001001). ’ REFERENCES (1) Kartha, S.; Grimes, P. Phys. Today 1994, 47, 54. (2) Murray, E. P.; Tsai, T.; Barnett, S. A. Nature 1999, 400, 649.

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