Ionic conductivities, sintering temperatures and microstructures of bulk ceramic CeO2 doped with Y2O3

Solid State Ionics 134 (2000) 89–102 www.elsevier.com / locate / ssi

Ionic conductivities, sintering temperatures and microstructures of bulk ceramic CeO 2 doped with Y 2 O 3 Chunyan Tian, Siu-Wai Chan* Materials Science and Engineering, School of Engineering and Applied Science, Columbia University, New York, NY 10027, USA

Abstract Ionic conductivities of CeO 2 :Y 2 O 3 bulk ceramic were investigated with different sintering temperatures and correlated with the resulted microstructure. Lower sintering temperatures (T # 14008C) were found to give much higher overall DC conductivities (e.g. sDC | 7 3 10 23 S / cm at 7008C for a 4% Y 2 O 3 -doped sample sintered at 14008C). The samples sintered at lower temperatures showed higher grain boundary conductivities than those sintered at the traditional sintering temperature, 15008C. The model involving non-resistive grain boundaries can be employed to explain the lower grain boundary resistivities in our samples of low sintering temperatures. These samples were examined by scanning transmission electron microscopy (STEM) with energy dispersive X-ray (EDX) and electron energy loss spectroscopy (EELS). Most of the boundaries ( . 90%) were found precipitate-free in the small grain samples, and a higher Y/ O ratio was observed at all these boundaries examined. The lower sintering temperatures suppress grain growth giving rise to small grain size (below 1 mm). The finer grain size provides large grain boundary areas for impurities to precipitate and solutes to segregate. Under such condition, there are insufficient impurities to form continuous precipitate layers at all boundaries, such as ion transport blocking layers at boundaries are not fully formed. At the same time, there are insufficient Y 9Ce ions for all the boundaries in the fine grain samples while the mobility of the Y 9Ce ions is low at low sintering temperatures to form well developed space charged regions at these boundaries to abate transboundary ionic transport. These three combined effects have abated some of most resistive mechanisms for ionic transport across boundaries.  2000 Elsevier Science B.V. All rights reserved. Keywords: CeO 2 doped with Y 2 O 3 ; Grain boundary conductivity; Ionic conductivity; Microstructure; Sintering temperature PACS: 61.16.Bg; 61.72.Mm; 66.10.Ed; 66.30.-h; 73.61.Tm; 81.05.Je; 81.10.Jt.

1. Introduction The electrical conductivity of ceria has been extensively investigated with respect to different dopants (Ca 21 , Y 31 , La 31 , Gd 31 , Sc 31 , Sr 21 , Sm 31 ) [1–3] and dopant concentration [4,5]. It was found that the electrical conductivity of doped ceria de*Corresponding author. Fax: 11-212-854-7081. E-mail address: [email protected] (S.-W. Chan).

pends on the type and the concentration of dopants. All of the doped ceria compositions showed an increase in the conductivity over undoped ceria, while those doped with Gd 31 , Sm 31 and Y 31 exhibited the highest conductivities [6–8]. Although all of these dopants have extensive solubility in the fluorite crystal structure of ceria, yttria (whose solubility in ceria is 20.5 mol% at 15008C) [9] was reported to be most soluble in the ceria lattice with excellent ionic conductivity. Comprehensive studies

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on Y 2 O 3 -doped CeO 2 have shown its conductivity varied substantially with the dopant concentration (0.05–40 mol% Y 2 O 3 ) [3,8]. Wang et al. [3] showed that the change of the lattice conductivity with dopant concentration corresponded to the change in the activation energy of oxygen vacancy diffusion. They found that a minimum in the activation energy appeared at a dopant concentration of | 4 mol% Y 2 O 3 , corresponding to the maximum lattice conductivity sl . A later observation by Sarkar and Nicholson gave similar results with maximum lattice conductivity at | 3 mol% Y 2 O 3 in CeO 2 [10]. A recent study by Balazs and Glass [8] showed that the maximum lattice conductivity of yttria-doped ceria occurred at a composition of Ce 1.84 Y 0.16 O 1.92 (8 mol% Y 2 O 3 ). There are differences among these lattice conductivity findings; however, no such systematic study of the grain boundary conductivity as related to the dopant concentration has been reported. In polycrystalline ceramic materials, grain boundaries often have a significant influence on overall conducting properties. The starting ceramic powders often contain silicate impurities inadvertently incorporated during powder processing [11]. These impurities tend to accumulate at grain boundaries during sintering and grain growth. The partial or complete blocking of charge carriers by these segregated impurities and siliceous precipitates leads to higher grain boundary resistivity [11,12]. The grain boundary resistivity is strongly influenced by impurities; it is small in relatively pure materials and increases with the impurity content (such as Si) appearing at grain boundaries [13]. Gerhardt et al. [12], using scanning transmission electron microscopy and transmission electron microscopy, found that the presence of Si in the form of a continuous glassy phase is responsible for the large decrease in the effective DC conductivity of ceria solid electrolyte. The contribution from the resistivity of ionic transport across grain boundaries can lower the overall DC conductivity as much as several orders of magnitude. Besides precipitates of extraneous phases at grain boundaries, other problems, such as imperfect alignment between grains (on an atomic scale), solute segregation and oxygen depletion resulting from elastic strain or space charge in the vicinity of the grain boundaries, can all increase the

grain boundary resistivity. To simplify further discussion, we define the precipitate related grain boundary resistivity as ‘extrinsic’ boundary resistivity and remnant boundary resistivity for precipitatefree intact boundaries as ‘intrinsic’ boundary resistivity. Earlier studies [13–15] reported that the grain boundary resistivity changed with the grain size. When the grain size decreases from a few microns to sub-micron level, the grain boundary showed unusually high conductivity ( s˜ ). No second phase at the grain boundaries was found, suggesting that the high grain boundary conductivity for samples of small grain size is related to the absence of a secondary phase at grain boundaries. The grain boundary resistivities have been proposed to result from different segregated elements [16] and the variation of grain boundary structure [17], but no confirmative experimental work was conducted. In our thin film work [18], we have found the lower substrate temperatures (700–9008C) of deposition contribute to higher boundary conductivities and an overall higher DC conductivity. In this work, we first showed the correspondence of the grain boundary conductivities with dopant concentrations and sintering temperatures. Then we studied the solute segregation and oxygen depletion effects on the grain boundary resistivities using microanalytical technique on the small-grain-size samples that showed high grain boundary conductivities.

2. Experimental

2.1. Pellet preparation The CeO 2 powders used to make ceramic pellets were from Alfa Aesar (99.99% pure), and Y 2 O 3 powders from Aldrich and Alfa Aesar (99.999% pure). The powders were carefully weighed, mechanically dry mixed in a desired proportion, then ground and cold die-pressed into pellets followed by uniaxial pressing at 20 000–40 000 p.s.i. without binder. The pellets were 2 cm in diameter, and the thickness ranged from 0.5 to 1 cm. The pellets, covered by appropriate powder to prevent contamination, were first sintered at 12008C for 48 h, then re-ground and pressed into pellets of the same

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dimension. The final sintering of these samples was conducted at various sintering temperatures between 8008C and 15008C in air for 4 days. The densities of the pellets were found to be of 68–95% of theoretical density.

91

Philips FEG CM-20 scanning transmission electron microscope (STEM). Both energy dispersive X-ray (Be window, EDAX) and electron energy loss (Model 666, Gatan) spectra were taken at 200 kV with a probe size of | 2 nm for 60 s.

2.2. Electrical measurements 3. Results and discussion The sintered pellet samples were cut and polished to 0.6–1 mm thick, and annealed at 9008C for 8–12 h. Silver paste was painted on both sides of the pellets with area of 0.4–1.5 cm 2 as electrodes. The samples with Pt wires attached to the electrodes were fired at 7008C to ensure a good bond between the sample surfaces and silver paste, then result in less electrode effect. The AC impedance was measured from 0.1 Hz to 10 MHz, in the temperature range between room temperature and 7008C, using a Solartron 1260 impedance analyzer. The measurement procedure was described in detail elsewhere [19].

2.3. Microstructural analysis To check for compositional uniformity of Ce and Y and impurity distribution in the pellets, energy dispersive analysis (EDS) by X-ray was used to examine the samples. The EDS was performed on a Cambridge SEM model 250Mk 2, with Kevex-8000 EDS attachment. The accelerating voltage, beam current, tilt angle, data collection time and spot size were all maintained the same for all samples examined by EDS. In order to observe the morphology of the pellets and measure the grain size, the samples were examined using SEM (Cambridge 250Mk 2). The surfaces of the pellets for microscopic observation were first polished to obtain a flat surface and followed by etching in a solution of hydrofluoric acid and nitric acid, then coated with carbon film to avoid surface charging when the SEM electron beam bombarded the surface. TEM specimens were prepared by mechanical polishing, dimpling and ion-milling to achieve electron transparency [19]. TEM observations were done on a JEOL E100 microscope with an accelerating voltage of 100 kV. Microanalysis on the 4% doped, 12008C sintering sample was conducted using a

3.1. Ionic conductivities The complex impedance plots were analyzed by fitting with different equivalent circuits [20], to obtain the resistance and capacitance of different arcs at different temperatures. The conductivities are calculated in two ways. The first is to convert the resistance to the conductivity by consideration of sample geometry as in Eq. (1) 1 t s 5]3] Ri S

(1)

where t is the separation of the electrodes (usually the sample thickness) and S is the electrode area. The series grain boundary conductivity calculated from Eq. (1) with the appropriate R gb is the macroscopic grain boundary conductivity sgb , as it is calculated with R gb and the macroscopic dimension of the sample (thickness / area) [20]. However, to study grain boundary effects, the macroscopic grain boundary conductivity sgb cannot give one insight since the grain boundary area parallel to current flow is much smaller than that of the lattice. Thus, one needs to include grain boundary layer thickness and grain size to calculate the microscopic series grain boundary conductivity s˜ gb , defined as d 1 t d s˜ gb 5 sgb ] 5 ] 3 ] 3 ] L R gb S L

(2)

where d is the grain boundary layer thickness, L is the grain size, and R gb is the resistance of grain boundary. The method of calculating conductivities using Eqs. (1) and (2) is referred to as the geometrical method. To obtain s˜ gb , one needs to know the grain boundary layer thickness d and grain size L, which can only be obtained by microscopy analysis. A quick way to obtain the conductivities is to use the ]resistance, ]capacitance, and the permittivities and

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Eqs. (3) and (4). This is referred to as the RC method. 1 t Rl 5 ] ] sl S S Cl 5 el e0 ] t el e0 sl 5 ]] R l Cl

(3a) (3b) (3)

1 t d R gb 5 ] ] ] s˜ gb S L

(4a)

S L Cgb 5 egb e0 ] ] t d

(4b)

For microscopic grain boundary conductivity:

egb e 0 s˜ gb 5 ]] R gb Cgb

(4)

boundary conductivities calculated with Eqs. (1), (3), and (4) are shown in Fig. 1 for a 6% doped pellet. For comparison, the lattice conductivity from Eq. (1) and the microscopic grain boundary conductivity estimated using Eq. (2) from literature results [22] are included. It can be seen that the lattice and the microscopic grain boundary conductivities are consistent with each other, respectively. The microscopic grain boundary conductivity is about 3 orders of magnitude lower than the macroscopic grain boundary conductivity.

3.1.1. Sintering at 15008 C Relevant to the operating temperature range for intermediate SOFCs, we measure the lattice and grain boundary conductivities at 500 and 7008C (see Fig. 2) for different dopant concentration pellets sintered at 15008C. Previously, researchers [1,3,8,12,15,17] only used sintering temperatures at or above 15008C. The high temperatures were to

where el and egb are the permittivity of the lattice and grain boundary, respectively. Hereafter, grain boundary conductivities for all the samples will be calculated using the RC method by assuming egb 5 el , i.e. they are all microscopic grain boundary conductivities. This assumption is reasonable here since the thickness of a grain boundary phase may easily vary along one boundary by more than the uncertainty introduced by estimating egb 5 e1 [21]. The lattice permittivity el can be obtained from the measured impedance results combined with the sample geometry. The conductivities are plotted as a function of temperature. Activation energies for conduction are obtained by plotting the conductivity data in the Arrhenius relation for thermally activated conduction, Eq. (5). Ea s T 5 s0 exp(2] ) kT

(5)

where Ea is the activation energy for conduction, T is absolute temperature, and s0 is a pre-exponential factor. The lattice conductivities calculated with Eqs. (1) and (3) are expected to be the same, while the grain boundary conductivities from Eq. (1) and (4) will be different. The difference will be determined by the ratio between the grain boundary layer thickness and the grain size. The lattice and grain

Fig. 1. Comparison of the lattice and grain boundary conductivities calculated using different methods for 6% Y 2 O 3 -doped CeO 2 pellets. Lattice conductivities are the same for these methods. The microscopic grain boundary conductivity calculated using the RC method is consistent with that from the literature [22] estimated by assuming the grain boundary thickness. They are both several orders of magnitude lower than the macroscopic one obtained using the geometric method.

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Table 1 Summary of activation energies, pre-exponential factors, and permittivities of the pellets sintered at 15008C with different dopant concentrations a Dopant conc., CO (mol%)

Fig. 2. Lattice and grain boundary conductivities at 500 and 7008C for the CeO 2 pellets sintered at 15008C with different dopant concentrations.

ensure the homogeneity of dopants. The highest lattice conductivity at 5008C belongs to the 6% doped samples, while 8% doped samples show the highest lattice conductivity at 7008C in agreement with Balazs and Glass [8]. The grain boundary conductivity increases with dopant concentration, reaching a maximum at 10% and then decreases as dopant concentration increases to 15%. No grain boundary conductivities were obtained for 20 and 40% doped samples. They show very low lattice conductivities instead. The decrease of the lattice conductivity at high dopant concentrations was attributed to the development of oxygen vacancy traps 9 V ??O )? defect associates. The activation [3] by (Y Ce energy and pre-exponential factors from our measurements are listed in Table 1. Wang et al.’s results [3], are also listed showing good agreements. The lattice dielectric constants (el ) are also shown in Table 1, which were used to calculate microscopic grain boundary conductivities ( s˜ gb ).

3.1.2. Sintering temperatures ( T #14008 C) The effect of sintering temperature on the conductivities of the pellets was investigated for two Y 2 O 3 doping concentrations: 0.58 and 4.0%. For 4.0% Y 2 O 3 -doped CeO 2 pellets, seemingly only one arc is observed in all the impedance plots of the pellets

Activation energy, Ea (eV)

Pre-exponential (s T ) 0 (SK / cm)

Dielectric constant, el

0.58

0.83 latt 1.31 gb 0.83 latt-w

2.872310 5 6.219310 3 1.584310 5

41.5

4.0

0.78 latt 1.07 gb 0.79 latt-w

1.605310 5 3.299310 2 6.306310 5

51.9

6.0

0.86 latt 1.03 gb 0.87 latt-w

2.384310 6 2.487310 3 3.161310 6

56.2

8.0

0.94 latt 1.02 gb 0.92 latt-w

1.394310 7 1.026310 4 2.511310 6

52.9

10.0

1.0 latt gb 1.14 0.98 latt-w

1.457310 7 5 1.057310 3.979310 6

53.3

15.0

1.13 latt 1.16 gb 1.10 latt-w 1.25 latt

2.295310 7 4.188310 4 9.994310 6 2.264310 7

48.1

1.41 latt 1.43 latt-w

5.981310 6 7.939310 6

52.4

20.0 40.0

58.5

a Latt, lattice conductivity; gb, grain boundary conductivity; latt-w, lattice conductivity from previous results (from Wang et al. [3]).

sintered below and at 14008C as shown in Fig. 3, while samples sintered at 15008C exhibit two arcs. The conductivity increases with sintering temperature for the samples sintered below and at 14008C, as shown in Fig. 4. This will be discussed later. The pellets sintered at 15008C show the highest lattice conductivity and lowest grain boundary conductivity. The activation energies and pre-exponential factors for those samples are listed in Table 2. Different phenomena were observed for 0.58% Y 2 O 3 -doped CeO 2 pellets sintered at different temperatures. Two arcs appear in the impedance plots for all the samples sintered at 10008C and above. As shown in Fig. 5, the lattice conductivity increases with sintering temperature. The grain boundary conductivity reaches a maximum at a sintering temperature of 12008C, and then decreases with

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Fig. 4. Overall lattice and grain boundary conductivities at 500 and 7008C for 4.0% doped pellets sintered at different temperatures. The pellets sintered at 14008C show the highest overall conductivity.

Table 2 Summary of activation energies and pre-exponential factors of 4.0% doped pellets sintered at different temperatures a

Fig. 3. Impedance spectroscopy plots for 4% doped pellets sintered at (a) 14008C and (b) 15008C. All samples sintered below and at 14008C showed seemingly one arc; while those sintered at 15008C exhibited two arcs.

Sintering temp., T s (8C)

Activation energy, Ea (eV)

Pre-exponential (s T ) 0 (SK / cm)

800 1000 1200 1300 1400 1500

1.06 o 0.98 o 0.92 o 0.92 o 0.85 o 0.78 latt 1.07 gb

9.279310 3 3.481310 4 1.091310 5 3.38310 5 1.792310 5 1.605310 5 3.299310 2

a

increasing sintering temperature. The difference between the lattice and grain boundary conductivities increases with sintering temperature. A summary of the activation energies and pre-exponential factors of these samples is given in Table 3.

3.2. Microstructure and microanalysis 3.2.1. Density The densities of the pellets sintered at different temperatures were measured relative to the theoretical density. The densities of the 4% doped pellets were consistently smaller than those of the 0.58% doped pellets. In general, the pellet density increases with sintering temperature as listed in Table 4.

o, overall sample conductivity; latt, lattice conductivity; gb, grain boundary conductivity.

3.2.2. Microstructure and microanalysis Pellets of different dopant concentrations were checked by energy dispersive spectroscopy (EDS). Several different points were measured for each sample to check the uniformity of the Y dopant distribution. The number of counts for Y and Ce obtained from EDS were converted into concentration by using the following equation. I ma /I m (A) 5 (ZAF )CA

(6)

where Ia and I(A) are the intensities from the

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Fig. 5. Lattice and grain boundary conductivities at 500 and 7008C for 0.58% doped pellets sintered at different temperatures. Those sintered at 12008C exhibit the highest grain boundary conductivity.

Table 3 Summary of activation energies and pre-exponential factors of 0.58% doped pellets sintered at different temperatures a Sintering temp., T s (8C)

Activation energy, Ea (eV)

Pre-exponential (s T ) 0 (SK / cm)

800 1000

0.94 0.73 latt 1.04 gb 0.77 latt 1.39 gb 0.80 latt 1.42 gb 0.81 latt 1.60 gb 0.83 latt 1.31 gb

2.809310 2 5.231310 2 3.41310 4 7.179310 3 1.208310 8 2.935310 4 2.533310 6 7.206310 4 5.497310 6 2.872310 5 6.219310 3

1200 1300 1400 1500 a

Latt, lattice conductivity; gb, grain boundary conductivity.

specimen and a pure element or a well-defined compound standard, respectively; CA is concentration in weight percentage; Z is the atomic number correction; A is the absorption correction, and F is the fluorescence correction [23]. The results, as summarized in Table 5, show good agreement to the starting compositions.

95

Grain sizes for the 0.58% doped samples sintered at 15008C ranged from 1 to 10 mm as shown in electron micrographs (Figs. 6 and 7). There are precipitates and an amorphous phase observed at grain boundary tri-junctions. However, for 4% doped pellets sintered at 15008C, and no precipitates or thick amorphous networks were observed along grain boundaries or at grain boundary tri-junctions using SEM and TEM. The average grain size of these samples is 10 mm (see Figs. 8 and 9). For lower sintering temperatures at 1200 and 14008C, TEM observation for 4% doped pellets (Figs. 10 and 11) shows the grain sizes are about 200–400 nm as compared to 10 mm for those sintered at 15008C. Therefore, extensive grain growth happened at 15008C but not at 14008C, which agrees with the results of the cation transport being the limiting step for grain growth [24]. The cations are not mobile until high temperatures around 14508C. In addition, no continuous, wetting amorphous layer was found in these samples sintered between 1200 and 14008C. For the samples sintered at 12008C, there are discontinuous amorphous layers along grain boundaries and precipitates at grain boundary tri-junctions (see Fig. 10). As compared to those precipitates in the samples sintered at 12008C, fewer lenticular precipitates were observed along grain boundaries of the samples sintered at 14008C (Fig. 11); this may explain the higher grain boundary conductivities of the samples sintered at 14008C. Table 6 gives a summary of the observed precipitate morphology and grain size for samples sintered at different temperatures and dopant concentrations. Both Si and Y segregation were found in those precipitates, but only Y segregation was found at grain boundary tri-junctions free of the precipitate in EDS spectra [19]. Among the 4% samples, those sintered at 14008C display a precipitation morphology which has the least blocking effect on conductivity. They also exhibit the highest grain boundary conductivity.

3.3. Grain boundary resistivity model To understand the fundamental origin of grain boundary resistivity, two basic models have emerged to describe the nature and location of the grain boundary phase in ceramics. For simplicity, the conducting mechanisms of oxygen-ion transport

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Table 4 Relative densities of pellets having different dopant concentrations sintered at various temperatures Density of sintered pellets / theoretical density of 7.13 g / cm 3 (%)

Sintering temp., T s (8C) 800 1200 1300 1400 1500

0.58% doped CeO 2

4.0% doped CeO 2

72.7 78.0 78.0 86.2 95.2

67.7 70.0 70.7 76.4 95.0

Table 5 Comparison of the starting Y 2 O 3 dopant concentration with that in the sintered pellets, which was obtained using energy dispersive X-ray spectroscopy Count ratio (Y/ Ce) (%)

Atomic concentration ratio (CY /CCe ) (%)

Mole concentration (CY 2 O 3 /CCeO 2 ) (%)

Starting concentration (CY 2 O 3 /CCeO 2 ) (%)

12.7 32.0 49.6 63.56

8.06 20.3 31.5 40.33

4.03 10.15 15.75 20.16

4.0 10.0 15.0 20.0

across these grain boundaries are represented by equivalent circuit models. 1. The parallel model was proposed by Schouler et al. [25] and adopted by Bernard [26]. The grain boundaries are partially blocked, and the ionic current across grain boundaries was considered to flow in two paths in this case. One path of high resistance partially covers the grain boundaries with ionic transport limited to 10 23 –10 22 of the overall conductivity. The other one is through the clean grain boundaries of grain-to-grain contact, and the resistivity is determined by the nature of the grain boundaries. Since the grain boundary resistivity from intimate grain-to-grain contact is much smaller than that from glassy phases, the overall grain boundary resistivity is determined mainly by the nature of the clean grain boundary region and not by the glassy phase. 2. Brick layer model, proposed by Van Dijk and Burggraaf [13], assumes that the ceramic samples made up of highly conducting grains surrounded by a continuous and uniform grain boundary layer. Ion transport across grain boundaries must take place through this grain boundary phase.

Furthermore, the ionic conductivity of the glassy phase is about three orders of magnitude lower than the intragrain conductivity. Gerhardt et al. [12] found that the presence of a continuous Si-rich glassy phase is responsible for the large decrease in the effective DC conductivities of ceria-based solid electrolytes. The contributions from intergrain and intragrain resistances to the total electrolyte resistance can be determined from the grain size and the separation of the electrodes. A series equivalent circuit can be used for representation of this model when grain size is smaller than the distance between the electrodes. Table 6 Summary of grain size and precipitate morphology for samples sintered at different temperatures Dopant concentration (mol%)

Sintering temp., T s (8C)

Grain size (mm)

Precipitate morphology

0.58 4.0

1500 1500

10 1–10

4.0 4.0

1400 1200

0.2–0.4 0.2–0.4

Coarse, thick, continuous Thin, maybe continuous Fewest precipitates Some precipitates

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Fig. 6. Scanning electron micrograph using backscattered electrons showing the surface of 0.58% doped pellets sintered at 15008C. The grain size varies from |1 to 10 mm, and there are precipitates at grain boundary tri-junctions and along grain boundaries.

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Fig. 9. Transmission electron bright field image of grain boundary tri-junctions of 4% doped pellets sintered at 15008C. No thick amorphous network is evident, but the grain boundary tri-junction and grain boundaries are not abrupt. Fig. 7. Transmission electron micrograph showing the thick grain boundary tri-junctions of 0.58% doped pellets sintered at 15008C, which is consistent with the observation using scanning electron microscopy as shown in Fig. 6.

The conductivities of all the pellets sintered at 15008C can be explained by adopting the brick layer model, i.e. the low grain boundary conductivities are due to the thin layer of glassy phase incorporated

Fig. 8. Scanning electron micrograph using backscattered electrons showing the surface of 4.0% doped pellets sintered at 15008C. No thick grain boundaries were observed.

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99

Fig. 10. Transmission electron bright field image of 4% doped pellets sintered at 12008C, showing grain size (|200–400 nm) and precipitates at grain boundary tri-junctions and along grain boundaries.

into grain boundaries during sintering. As the dopant concentration increases, the oxygen vacancy concentration will increase accordingly; thus, the grain boundary conductivity increases with dopant concentration. However, when the dopant concentration reaches some point, the oxygen vacancy V ??O and the 9 tend to associate and form a [V O?? Y Ce 9 ]? dopant Y Ce ?? pair, resulting in a decrease in [V O ] [3]. Furthermore, when the dopant concentration is high enough, the dopants may tend to segregate at the grain boundaries and form trapping layers for oxygen ion transport. However, full segregation would not be reached at low sintering temperature because of the 9 . low mobility of Y Ce Based on the microstructure and microanalysis results given earlier, the parallel model may give a better explanation of the conductivity behavior of those small-grain-size samples, which exhibit much higher grain boundary conductivity than those sintered at 15008C. The grain boundary area is so large that the finite amount of impurity contained in these samples is not sufficient to form a continuous and uniform glassy phase layer along grain boundaries. On the other hand, the sintering temperature may not

be high enough for the impurities to form a continuous and uniform glassy phase layer along grain boundaries, so they segregate discontinuously at the grain boundaries leaving the remaining grain boundary areas with clean grain-to-grain contact. Therefore, the transport of oxygen ions is through two paths, and the overall grain boundary conductivity is mainly determined by the clean boundaries with ‘intrinsic’ grain boundary conductivity. For the 0.58 and 4.0% doped samples sintered at low temperatures, the grain boundary conductivity increases with the dopant concentration. For 4.0% doped samples sintered at low temperatures, the grain boundary conductivities are comparable with their lattice conductivities such that the two arcs cannot be separated as distinct semicircles in complex impedance plots. Hence, they show very high overall DC conductivities as compared to those sintered at 15008C. Grain Boundaries in these samples show an absence of a continuous layer. Samples with 4% dopant yield highest DC conductivity when sintered at 14008C. Samples with 0.58% dopant yield highest DC conductivity when sintered at 12008C. However,

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even with these samples two semicircles were observed in impedance spectroscopy plots. Here the 0.58% dopant concentration maybe not high enough to decrease the grain boundary conductivity to be comparable to that of the lattice. The semicircles due to the lattice and grain boundaries can still be separated. This is different from the single semicircle observed for the 4% dopant samples sintered at 14008C and below.

3.4. Space charge model at grain boundaries Since most of the grain boundaries in this study were precipitate-free, we will discuss in the following what contribute to the ‘intrinsic’ boundary resistivity. In particular, we will examine the resistive contribution from space charge layers at grain boundaries. Frenkel first proposed the existence of a space charge region near lattice discontinuities (free surfaces and grain boundaries) in an ionic solid [27]. Grain boundaries have been assumed to act as infinite sources and sinks for vacancies. At thermodynamic equilibrium the grain boundaries of an ionic crystal may carry an electric potential resulting from the accumulation of the charged point defects. The charge excess is a consequence of the different free energies of formation of cation and anion vacancies for pure ionic crystals [28]. This core potential is compensated by a space charge region of the opposite sign adjacent to the boundary. In semiconducting samples, electronic carriers (holes or electrons) respond to the core potential and are depleted or accumulated forming space charged regions. When electronic carriers are unavailable as in ionic materials, small concentrations of aliovalent ions (from dopants or impurities) which possess a charge different from that of the host ions of the crystal, could diffuse in response to the electrostatic potential at grain boundaries to form segregated or depleted layers and give rise to a space-charge region at grain boundaries. For our samples, Y 2 O 3 -doped CeO 2 pellets, we propose that the effectively positively charged oxygen vacancies are the predominant defects at grain boundary cores. Since oxygen anions are larger and they will run into tight spots at most boundaries. It will be energetically favorable to leave those spaces unfilled (i.e. oxygen vacancies) than clamping in oxygen anions as proved by computer simulations of

oxide boundaries [29]. The center core then will have a positive potential because of excess oxygen vacancies. In order to maintain long-range charge 9 segregates to neutrality, the negatively charged Y Ce the grain boundaries to form space charge layers, as schematically shown in Fig. 12. The formation of this boundary charge requires a redistribution of the ions; effectively negatively charged Y 9Ce segregates at the space charge layer, while positively charged V ??O accumulates at grain boundaries. Therefore, both Y segregation and oxygen depletion are expected at grain boundaries. What happens to an oxygen vacancy as it crosses a boundary with well developed positive potential 9 core-layer sandwiched by negatively charged Y Ce segregated space-charge layers? It has a higher probability to be trapped to form an associate defect 9 which has a higher concentration in the with Y Ce space charged layers. This will contributed to the boundary resistance even when amorphous siliceous layers are absent. In order to obtain the composition profile across grain boundaries, we perform a line scan in STEM [30]. The probe was scanned across grain boundaries at a step of a few nanometers and a spectrum was collected at each point with EDS and EELS simultaneously. This type of profile is effective for revealing the presence of segregation or depletion, and can be conducted with a spatial resolution of 5 nm with FEG CM-20 used in our work. The STEM results shown in Fig. 13(a–c) are consistent with the space charge model proposed above. We estimated the Debye length of the space charge layer to be 26 ˚ in a 4% doped sample sintered at 12008C. This A estimate agrees with the microanalysis results of 5–10 nm as shown in Fig. 13(b,c) considering a 5-nm spatial resolution. It also can explain that the higher boundary conductivity for samples sintered at lower temperatures than the traditional 15008C. In these samples with lower sintering temperatures, the space charge layers are not well-developed to significantly decrease the oxygen vacancy transport across boundaries. The lower sintering temperatures help to lessen the space-charge formation because higher temperatures are needed to form space9 mobility in CeO 2 can be the charged layers. The Y Ce limiting step of forming well-developed spacecharged layers. Previous results on grain size in CeO 2 have given strong evidence to support the

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Fig. 11. Transmission electron bright field image of 4% doped pellets sintered at 14008C. Only a few lenticular precipitates are visible along grain boundaries.

mechanism of cation transport being the rate-limiting step in grain growth upon different sintering temperatures [24].

4. Conclusions The sintering temperature has a significant effect on the pellet conductivities. We find that by going to

Fig. 13. STEM image (a) of the grain boundary for EELS analyses and the concentration profiles of Y and O cross grain boundary of the 4% doped pellet sintered at 12008C. The changes of the concentration inside grain and at grain boundary show Y segregation (b) and O depletion (c) at the grain boundary. Fig. 12. Schematic diagram of a grain boundary in yttria-doped ceria showing excess oxygen vacancies at the grain boundary core 9 segregation. and adjust space-charge layers with Y Ce

a lower sintering temperature below the traditional 15008C, we get a higher overall DC conductivity. The highest lattice conductivity occurs for the dopant

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concentration of 4–8% at different measuring temperatures, while maximum grain boundary conductivity was observed at 10% dopant concentration for samples sintered at traditional sintering temperature of 15008C. The conductivities are more determined by pre-exponential factors than activation energies as temperature increases. The highest overall DC conductivity occurs at 1400 and 12008C sintering temperatures for 4 and 0.58% doped samples, respectively. Our results support that the overall DC conductivities for solid electrolytes can be improved dramatically by choosing an appropriate sintering temperature. The high grain boundary conductivities for lower temperature sintered samples are due to: (1) lower impurity concentration in grain boundaries, resulting from small grain size and larger grain boundary areas, with few precipitates and discontinuous blocking layers; (2) lower segregation of the solutes and less well developed space charge regions at grain boundaries partly because of small grain size and low mobility of cations at lower temperatures, respectively. The activation energies for overall sample conductivities are about 0.1–0.2 eV higher than those of lattice conductivities for 4% doped samples sintered at low temperatures. Similar results in CeO 2 :Y 2 O 3 thin films with higher overall DC conductivities were observed [19,31–33]. These films were processed at low deposition temperatures.

Acknowledgements We thank Prof. A. S. Nowick for helpful discussions. The support from National Science Foundation under grant DMR-93-50464 is appreciated.

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