Probing Charge-State Distribution at Grain Boundaries Varied with

Article pubs.acs.org/JPCC

Probing Charge-State Distribution at Grain Boundaries Varied with Dopant Concentration for Ceria Ceramics Kiminori Sato* Department of Environmental Sciences, Tokyo Gakugei University, 4-1-1 Koganei, Tokyo 184-8501, Japan ABSTRACT: Knowledge of dopant-concentration dependent grain-boundary conductivity is a prerequisite for designing sophisticated nanostructured electrolytes as oxide ceramics but is not yet fully understood for the lack of quantitative information on the charge-state distribution there. Here, Gd-doped ceria (GDC) with nano- and micrometer sized grains, referred to as nano- and micro-GDC, is studied as a function of dopant concentration by means of element-specific positron annihilation spectroscopy. It is demonstrated that positrons are able to probe the negatively charged spaces adjacent to the grain boundary core in addition to the electrons bound to Ce atoms in the bulk. The atomic concentration of negatively charged space is ∼10−7 both for micro- and nano-GDC and increases as the Gd dopant increases. In contrast to that, the size of negatively charged space shrinks with the dopant concentration due to a charge effect arising from the electrons associated with the Gd dopant. The volume occupancy of negatively charged space satisfactorily explains the grainboundary conductivity dependent on the dopant concentration as well as the grain size in terms of the space-charge theory.

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core.6,8,10−15 The presence of an excess positive charge in the core is generally explained in terms of an excess of structurally necessary positively charged oxygen vacancies. The width of space charge zone λ depends on the concentration of the majority carrier in the bulk cmaj,∞ as shown by the following expression6,8

INTRODUCTION Rare-earth-doped ceria is one of the best solid electrolytes for environmentally friendly solid-state fuel cells owing to its high oxygen vacancy conductivity at intermediate temperatures.1 A well-known example is Gd-doped ceria (GDC) with the composition of Ce0.8Gd0.2O2−δ, which is an excellent ceramic as an oxygen ionic conductor due to its high conductivity and structural stability.2 In terms of Kröger−Vink notation3 which is extensively used in defect chemistry, the doping reaction of GDC is described as Gd 2O3 → 2Gd′Ce +

3OOX

+

V •• O

λ=

2 2 2zmaj e cmaj, ∞

(2)

where ε0, εr, kB, T, zmaj,∞, and e are the permittivity of free space, the relative dielectric constant, the Boltzmann constant, the absolute temperature, the charge number of majority carrier, and the elementary charge, respectively. Under the condition of uniform dopant concentration, the Mott−Schottky approximation16 can be applied, yielding the width of the spacecharge layer defined as

(1)

generating the negatively charged Gd dopant on the Ce4+ lattice site (Gd′Ce), oxide ion (OXO), and positively charged oxygen vacancy (V··O). The ionic conductivity in GDC originates from the migration of oxygen ions4,5 resulting from the presence of negatively charged acceptor defects (Gd′Ce). Such a transport of ions inside the grains and across grain boundaries is generally known as bulk and specific grain boundary conductivities, respectively. Based on the results of electrochemical impedance spectroscopy, the grain-boundary conductivity for solid electrolytes has been found to be lower than that of the bulk by several orders of magnitude.6,7 The highly resistive ionic transport at the grain boundary could stem from the segregation of undesired impurities there for the imperfectly synthesized ceria samples, which has been identified by high-resolution transmission electron microscopy (TEM) observations.8,9 On the one hand, ionic transport has been found to degrade at the grain boundaries for ceria, since a depletion of positively charged oxygen vacancies occurs in the negative space-charge layers adjacent to the positively charged grain-boundary © XXXX American Chemical Society

ε0εrkBT

λ=

2ε0εr Δφ0 zdopecdop, ∞

, (3)

where Δφ0, zdop, and cdop,∞ are the space charge potential resulting from the excess positive charge present in the core of the grain boundary, the charge number of dopant, and the concentration of dopant in the bulk, respectively. It is evident from these equations that the width of the space-charge layer decreases with the square root of the dopant concentration for the given space-charge potential. An electrochemical approach Received: August 11, 2017 Revised: August 30, 2017 Published: August 31, 2017 A

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+2.5) × 10−3 m0c and the core electron area ± (25 to 35) × 10−3 m0c with normalization to the total area of the spectrum, respectively.

has suggested that the specific grain-boundary conductivity increases with increasing dopant concentration.6,15,17,18 The facts mentioned above imply that there exist two kinds of negatively charged defects for polycrystalline ceria ceramics. One is negatively charged point defects such as acceptors (e.g., GdCe ′ ) or reduced cerium ions (Ce3+ namely CeCe ′ ) homogeneously distributed in the bulk. Another one is an accumulation of defects in the proximity of grain boundaries leading to the formation of two-dimensional negatively charged space region. In this contribution, we intend to probe such negatively charged defects by making full use of a positron, which is a positively charged particle with mass identical to that of an electron. A feature of positron as antimatter enables us to probe the negatively charged spaces locally present in materials, being attracted through Coulombic interaction with them. Exemplary representatives are the negatively charged part localized on polar elements in polymers19,20 as well as the metallic lattice vacancies21,22 that are negatively charged relevant to the matrix due to missing of positive ion core. Our previous work has demonstrated that the negatively charged spaces on an angstrom-scale adjacent to the grain boundary core for GDC and Li−-GDC can be directly probed by positrons.23 This is on the basis of positron trapping at the negatively charged space, providing the space size and its chemical surrounding through the measurements of positron lifetimes by positron lifetime spectroscopy and momentum distributions of annihilation photons by coincident Doppler broadening spectroscopy, respectively. In this study, the combination of the two spectroscopic techniques, called element-specific positron annihilation spectroscopy, was employed as the direct probe of negative spaces in GDC. Aside from the electrochemical studies exclusively conducted so far, we focus on the concentration of negatively charged space and its size varied with the dopant concentration.

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RESULTS AND DISCUSSION Positron lifetime spectroscopy yields two components of positron lifetimes τ1 and τ2 for micro- and nano-GDC regardless of the Gd concentration. In Figure 1, positron

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EXPERIMENTAL SECTION The GDC samples with the nano- and micrometer-sized grains containing the Gd content of 0, 0.15, 1, and 6 atom % were prepared following the procedure described in the literature.24 Thereafter, we call the samples with the nano- and micrometersized grains nano- and micro-GDC, respectively. The nano- and micro-GDC samples with the Gd content of 10 atom % were obtained from the nanocrystalline powder commercially available from Sigma-Aldrich under the same pressing and sintering conditions. This raw nanocrystalline powder possesses a larger grain size than those of nano- and micro-GDC prepared at 0, 0.15, 1, and 6 atom %. Positron lifetime and coincident Doppler broadening (CDB) spectroscopy were performed at room temperature. The positron source (22Na), sealed in a thin foil of Kapton, was mounted in a sample-source-sample sandwich for the measurements. Positron lifetime spectra (∼1 × 106 coincidence counts) were recorded with a digital oscilloscope-based system in which the time resolution of 200 ps full-width at half-maximum (fwhm) was achieved.25,26 The lifetime spectra were numerically analyzed using the POSITRONFIT code.27 In CDB spectroscopy, the energies of the two annihilation quanta E1 and E2 were measured with a collinear setup of two high-purity Ge detectors. The spectra were obtained by cutting the E1, E2 spectra along the energy conservation line E1 + E2 = (1022 ± 1) keV, taking into account the annihilation events within a strip of ±1.6 keV. CDB spectra were analyzed by taking the S and W parameters covering the valence electron area (−2.5 to

Figure 1. Positron lifetimes τ1 (a) and τ2 (b) together with the relative intensity I2 (c) of τ2 for micro- (black open symbols) and nano-GDC (red solid symbols). The solid lines are drawn for guiding the eye.

lifetimes τ1 (a) and τ2 (b) are plotted as a function of Gd content together with the relative intensity I2 (c) of τ2. The values of shorter lifetime τ1 obtained for micro- and nano-CeO2 corresponding to the Gd content of 0 atom % are 178 and 182 ps, respectively. They are close to the lifetime of YSZ single crystal (see dashed line in (a)),23 which thus correspond to annihilation in the bulk. The positron lifetime τ1 for nano-CeO2 is slightly longer than that of micro one could arise from a little localized positron state due to structural imperfections such as local distortion caused by the suppression of CeO2 grains. This is also anticipated from the relative intensity I2 for nano-CeO2 that is higher than that of the micro one (see Figure 1c). The longer positron lifetimes τ2 of ∼300 ps and ∼325 ps obtained for the micro- and nano-CeO2 samples, respectively, are likely to be resultant from the positron-localized state such as positron trapping into a certain space (see Figure 1b). The analysis of electrochemical impedance data with the space charge theory has suggested that the grain-boundary core for B

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the grain boundaries is more easily completed for micro-GDC than that of nano-GDC. This causes saturation in lowering the positive space charge potential of grain boundary core at the Gd content higher than 1 atom %. Thus, Gd doping is less effective in the shrinkage of negatively charged spaces adjacent to the grain boundary core. In addition to that, there could be an effect of Gd3+ segregation at the grain boundaries for nanoGDC as often reported by TEM observations,8,9 which will be actually confirmed by the element-specific data of Figure 2.

nano-CeO2 is positively charged due to enrichment of oxygen vacancies there.11 The positively charged grain boundary core could create the adjacent negatively charged spaces, acting as a positron trapping center.23 The grain boundary for nano-CeO2 is expected to be loosely packed with oxygen vacancies more than that of micro-GDC, by which the negatively charged spaces get larger. This could be the reason why the positron lifetime τ2 for nano-CeO2 is higher than that of micro one. It is emphasized that the negatively charged spaces directly identified here are the origin of space-charge layer, which has been considered in the context of space charge theory. The positron lifetime τ1 for micro- and nano-GDC monotonously decrease as the Gd content increases (see Figure 1a). It is expected that the negatively charged Gd dopants (GdCe ′ ) and positive oxygen vacancies (VO·· ) are associated with each other owing to the Coulombic interaction, yielding negative charges due to an inhomogeneity of charge distribution in the bulk. Extended X-ray absorption fine structure (EXAFS) studies have suggested the local structural distortion near cation site28 as well as the formation of defect associates between the dopant and oxygen vacancy.29 The positively charged oxygen vacancies could not attract a positron, positively charged electron. In the light of the fact that positrons preferentially interact with the negative charges in materials, the positron lifetime τ1 is thus attributable to the annihilation with the electrons in the bulk. With increasing the Gd content more negative charges are available in the matrix, leading to an increase in the density of negative charges. This is well seen in the monotonous decrease of positron lifetime τ1 together with the Gd content. The consistently longer positron lifetimes τ1 with the higher relative intensities I2 obtained for nano-GDC could be caused by positron localization similarly to nano-CeO2. The positron lifetimes τ2 for micro- and nano-GDC are consistently higher than ∼240 ps in the whole Gd content range, which are ascribed to annihilation within the negatively charged space since the positively charged grain-boundary core arising from the oxygen vacancies is present for GDC as well. They are in the range of 275−320 ps at 1 atom % and decrease down to ∼250 ps with increasing the Gd content to 10%, indicating the shrinkage of negatively charged space from the vacancy-cluster sized to monovacancy-sized one,23,30,31 that is roughly by ∼1 Å in radius. The positron lifetime τ2 for microGDC decreases drastically with the Gd content to 1 atom % and then slightly decreases thereafter, whereas the gradual decrease with the Gd content is observed in the whole Gd content range for nano-GDC. In the proximity of grain boundaries, the negative charge resulting from Gd doping is expected to be localized more than those in the bulk owing to the structural disorder without periodicity. The number of electrons increases with increasing the Gd content, enriching the electron concentration in the vicinity of grain boundaries. This causes the reduction in the size of negatively charged spaces adjacent to the positive grain boundary core as can be seen in Figure 1b. It is of interest that the shrinking tendency of negatively charged spaces with the Gd dopant differs between micro- and nano-GDC (see Figure 1b). Let us first discuss on the data of micro-GDC, where such a shrinking is saturated for a Gd concentration >1 atom %. The grain boundary structure for micro-GDC is expected to be denser with less oxygen vacancies than that of nano-GDC, the core being less positively charged. The above-mentioned enrichment of electron concentration at

Figure 2. (a) Correlation between S and W parameters taken from coincidence Doppler broadening spectra for micro- (blue squares) and nano-GDC (red triangles). The data for GDC sintered at the temperatures from 773 to 1573 K (open circles) together with those for Gd2O3 (black solid circles) and CeO2 (blue solid square and red solid triangle) are presented for comparison. The solid and dashed lines with the data of Gd2O3 and CeO2 are drawn for guiding the eye, respectively. (b) Blown-up section of nano-GDC. The arrows are drawn for guiding the eye.

When Gd interstitially segregates into the core, the positive space charge potential of grain boundary is enhanced, by which the lowering of positive space charge potential is lessened. The shrinkage of negatively charged space thus occurs more slowly with the Gd content for nano-GDC than that of micro one as seen in Figure 1b. The relative intensity I 2 of the positron lifetime τ 2 corresponding to annihilation at the negatively charged space is lower than ∼30% in the whole Gd content range for microGDC (see Figure 1c). The lower intensities indicate that positrons are dominantly annihilated in the bulk for microGDC. The relative intensities I2 for nano-GDC are consistently higher than those of micro-GDC, indicating an availability of negatively charged spaces more than those in micro-GDC. This is easily understandable with respect to the difference of grain boundary structure between micro- and nano-GDC. The grain boundaries for nano-GDC could be concentrated with the structure being loosely arranged than that of micro-GDC. The concentration of oxygen vacancies for nano-GDC is resultantly higher than that of micro-GDC at the grain boundaries, leading to the increase in concentration of adjacent negative charges spaces. The intensity I2 for micro-GDC increases from ∼10 to ∼30% as the Gd content increases, demonstrating that the negatively charged space increases along with the Gd dopant. The intensity I2 for nano-GDC increases up to ∼90% with increasing the Gd content to 6 atom % and then decreases down to 74% at 10 atom %. The increases of I2 commonly observed for micro- and nano-GDC are resultant from the increases of positive oxygen vacancies at the grain boundary C

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The Journal of Physical Chemistry C core caused by Gd doping, where negative spaces are increasingly formed with the Gd dopant. The decrease of I2 seen at 10 atom % could be attributable to the raw materials with larger grain sizes employed for the present synthesis (see the Experimental Section). This reduces the concentration of grain boundaries suppressing the formation of negatively charged spaces therein. Figure 2a shows the correlation between S and W parameters (S−W correlation) for micro- (solid and open squares) and nano-GDC (solid and open triangles) together with that of Gd2O3 (solid circle). The data of GDC at the Gd content of 0 atom % correspond to those of CeO2, which are denoted as the symbols of solid square and triangle for micro- and nano-CeO2, respectively. Here, the present data are compared with the correlation of GDC sintered at the temperatures from 773 to 1573 K (open circles) taken from our former work,32 for the purpose of discussing the variation of elemental environment around the positron annihilation site. Upon sintering-induced densification, the valence-electron contribution to the momentum distribution of annihilation photon gets lower due to the shrinkage of open spaces, whereas the contribution of core electron having elemental information around positron annihilation site gets higher in the same degree. Thus, the S parameter linearly varies with the W parameter when the chemical environment around positron annihilation site is unchanged during sintering.33 The data in the correlation are located on the straight solid line of Gd2O3 (Gd2O3 line) in the temperature range up to 973 K, indicating that the surrounding atoms of positron annihilation sites are similar to those of Gd2O3. This element-specific result captures the picture that positrons probe the electrons of Gd residing in the grain boundaries of GDC in the sintering temperature up to 973 K due to insufficient densification with the relative density less than ∼55%.34 The data deviate from the Gd2O3 line and are in turn located on the straight dashed line of CeO2 (CeO2 line) above 1073 K, as the chemical surroundings are getting toward to those of CeO2. Basically, the data of micro- and nano-GDC are located on the CeO2 line regardless of Gd content, demonstrating that the positron annihilation sites are dominantly surrounded by Ce atoms. The S-W correlation for the nano-GDC are sensitively influenced by Gd doping, though the correlation stays on the Ce line for micro-GDC. In Figure 2 the blown-up section of nano-GDC is shown as the inset (b), in which the correlation of nano-GDC at the Gd content of 1 atom % is significantly deviated from the CeO2 line as indicated by an arrow. The tendency of deviation becomes increasingly prominent when the Gd content is raised up to 6 and 10 atom % (see arrows). The downward deviation from the CeO2 to Gd2O3 line is the consequence of positron annihilation partially with the electrons of Gd atom. In the present positron lifetime spectroscopy, the bulk component is dominantly detected for micro-GDC, whereas the grain boundary is dominant component for nano-GDC (see Figure 1c). The elemental information deduced from the S-W correlation thus arises from the bulk and the grain boundary for micro- and nano-GDC, respectively. It is presumed that Gd3+ segregation develops at the grain boundaries for nano-GDC with increasing the Gd content, as it is supposed in the discussion on the shrinkage of negatively charged spaces above. For a more quantitative discussion of the negatively charged spaces identified here, the two-state positron trapping model34

was applied to the present positron annihilation data. The positron trapping rate κd is obtained from the following relation κd =

I τ + I2τ1 ⎞ 1⎛ ⎟ ⎜1 − 1 2 τ1 ⎝ τ2 ⎠

(4)

where τ1, τ2, I1, and I2 are the positron lifetime at the bulk, the positron lifetime in the negatively charged space, the relative intensity of τ1, and the relative intensity of τ2, respectively. Figure 3 shows the positron trapping rate κd for micro- and

Figure 3. Positron trapping rate κd for micro- (black open squares) and nano-GDC (red solid squares) plotted as a function of Gd content. The solid lines are drawn for guiding the eye.

nano-GDC plotted as a function of Gd content. The value of κd for micro-GDC gradually increases with increasing the Gd content, whereas nano-GDC exhibits the rapid increase in κd more than that of micro one with increasing the Gd content to 6 atom % and then the lower value at the Gd content of 10 atom %. The lowering of κd at 10 atom % for micro-GDC is resultant from the suppression of I2 caused by the raw material with larger grain size (see Figure 1c). The increases in κd for both micro- and nano-GDCs indicate that positrons tend to trap into the negatively charged spaces as the Gd content increases. As mentioned above, the grain boundary core for nano-GDC is looser with positive oxygen vacancies than that of micro-GDC. This results in the formation of negatively charged spaces adjacent to the grain boundary core more than those of micro-GDC (see Figure 1c), thus leading to the rapid increase in κd for nano-GDC. The atomic concentration of negatively charged spaces Cd can be estimated by using the relation κd = μd Cd

(5)

where μd is the specific positron trapping coefficient. Here, the constant specific positron trapping coefficient μd ∼ 2 × 1015 s−1 typical for negatively charged defect in ceramics35 was employed, since the value of μd could insignificantly decrease within the size reduction of ∼1 Å. Figure 4 shows the concentration of negatively charged space Cd as a function of Gd content. The Gd-content variations of Cd for micro- and nano-GDC exhibit a similar tendency to that of positron trapping rate κd in Figure 3. This is understandable since the amount of negative spaces efficiently contributes to the positron trapping rate as discussed above. In contrast to the fact that the negatively charged space shrinks as the Gd content increases, the concentration Cd increases both for micro- and nano-GDC. We thus introduced D

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the grain boundary core in addition to the electrons bound to Ce atoms in the bulk. The atomic concentration of negatively charged space estimated by the two-state positron trapping model is ∼10−7 both for micro- and nano-GDC and increases as Gd dopant increases. The increase in the concentration of negatively charged space for nano-GDC is more sensitive to the Gd dopant than that of micro-GDC owing to more oxygen vacancies available in the grain boundaries. Though the concentration of negatively charged space increases with the Gd dopant, the size shrinks both for micro- and nano-GDC due to the charge effect arising from electrons around the Gd dopant. The shrinking tendency is insensitive to the Gd dopant for nano-GDC due to Gd3+ segregation in the grain boundaries. The negatively charged spaces could correspond to the formation of negative-space-charge layer, in which the positively charged oxygen vacancies are depleted thus being less contribution to the grain-boundary ionic conductivity. The volume occupancy of negatively charged space introduced for the purpose of quantitative discussion on the spatial distribution of negatively charges satisfactorily explains the grain-boundary conductivity dependent on the Gd-dopant concentration as well as the grain size in terms of the space charge theory. The present findings are of particular importance not only for better understanding of ionic transport at the grain boundaries but also the design of sophisticated nanostructured electrolytes as oxide ceramics.

Figure 4. Concentration of negatively charged space Cd for micro(black open squares) and nano-GDC (red solid squares) plotted as a function of Gd content. The solid lines are drawn for guiding the eye.

the volume occupancy of negatively charged space ϵ through a simple equation 4 ϵ = πrd3cd (6) 3 where rd is the size of negatively charged space. The volume occupancy ϵ provides a good indication for the spatial distribution of negatively charged spaces, taking the shrinkage of negative space and the increase of concentration into account. Assuming the size reduction of ∼1 Å, eq 6 gives the decreases of ∼3.0 × 10−6 and ∼1.3 × 10−5 in the volume occupancy ϵ for micro- and nano-GDC, respectively, when the dopant concentration is raised up to 10 atom %. It is noted here that the decrease in volume occupancy ϵ for nano-GDC is an order of magnitude higher than that of micro-GDC. This demonstrates that the negatively charged spaces more efficiently disappear with the Gd content for nano-GDC than that of micro-GDC. According to the electrochemical studies of GDC by AvillaParedes et al.,6 the grain boundary conductivity increases as the Gd content increases and as the average grain size decreases. The volume occupancy of negatively charged spaces introduced here well explains the grain-boundary conductivity dependent on the dopant-concentration and grain size. The decrease of volume occupancy with the Gd dopant demonstrated above reduces the net width of negative-space-charge layer, in which the positively charged oxygen vacancies are less depleted. This contribute to an improvement of grain-boundary conductivity together with the increase in concentration of oxygen vacancies. Interestingly, the decrease of grain boundary width with Gd content was pointed out based on the results of space-charge analysis for electrochemical impedance data.5 For nano-GDC, the disappearance of negatively charged spaces efficiently occurs with the Gd content together with an enrichment of oxygen vacancies due to the loose grain boundary structure. Thus, the grain boundary conductivity is improved as the average grain size decreases, though the shrinkage process of negatively charged spaces is suppressed by Gd3+ segregation.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Kiminori Sato: 0000-0002-1622-5293 Notes

The author declares no competing financial interest.

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ACKNOWLEDGMENTS The author is indebted to Dr. Giuliano Gregori (Max-PlanckInstitut für Festkörperforschung) for supplying the present GDC samples and thorough reading the manuscript. This work was partially supported by a Grant-in-Aid from the Japanese Ministry of Education, Science, Sports and Culture (Grant No. 16K05394).

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REFERENCES

(1) Kilner, J. A.; Waters, C. D. The effects of dopant cation-oxygen vacancy complexes on the anion transport properties of nonstoichiometric fluorite oxides. Solid State Ionics 1982, 6, 253−259. (2) Taniguchi, T.; Watanabe, T.; Sugiyama, N.; Subramani, A. K.; Wagata, H.; Matsushita, N.; Yoshimura, M. Identifying defects in ceriabased nanocrystals by UV resonance Raman spectroscopy. J. Phys. Chem. C 2009, 113, 19789−19793. (3) Kröger, F. A.; Vink, H. Relations between the concentrations of imperfections in crystalline solids. Solid State Phys. 1956, 3, 307−435. (4) Ismail, A.; Giorgi, J. B.; Woo, T. K. On the atomistic interactions that direct ion conductivity and defect segregation in the bulk and surface of samarium-doped Ceria: A Genetic Algorithm Study. J. Phys. Chem. C 2012, 116, 704−713. (5) Yashima, M.; Takizawa, T. Atomic displacement parameters of ceria doped with rare-earth oxide Ce0.8R0.2O1.9 (R = La, Nd, Sm, Gd, Y, and Yb) and correlation with oxide-ion conductivity. J. Phys. Chem. C 2010, 114, 2385−2392. (6) Avila-Paredes, H. J.; Choi, K.; Chen, C. T.; Kim, S. Dopantconcentration dependence of grain-boundary conductivity in ceria: A space-charge analysis. J. Mater. Chem. 2009, 19, 4837−4842.

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CONCLUSIONS Gd-doped ceria (GDC) with the nano- and micrometer sized grains, referred to as nano- and micro-GDCs, was studied as a function of dopant concentration ranging from 0 to 10 atom % by element-specific positron annihilation spectroscopy. We succeeded in probing the negatively charged spaces adjacent to E

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Article

The Journal of Physical Chemistry C (7) Crumlin, E. J.; Mutoro, E.; Ahn, S. J.; Jose la O’, G.; Leonard, D. N.; Borisevich, A.; Biegalski, M. D.; Christen, H. M.; Shao-Horn, Y. Oxygen reduction kinetics enhancement on a heterostructured oxide surface for solid oxide fuel cells. J. Phys. Chem. Lett. 2010, 1, 3149− 3155. (8) Guo, X.; Sigle, W.; Fleig, J.; Maier, J. Role of space charge in the grain boundary blocking effect in doped zirconia. Solid State Ionics 2002, 154−155, 555−561. (9) Li, Z. P.; Mori, T.; Auchterlonie, G. J.; Zou, J.; Drennan, J. Direct evidence of dopant segregation in Gd-doped ceria. Appl. Phys. Lett. 2011, 98, 0931041−0931043. (10) Maier, J. Space charge regions in solid two-phase systems and their conduction contribution - I. Conductance enhancement in the system ionic conductor-‘Inert’ phase and application on AgCl: Al2O3 and AgC1: SiO2. J. Phys. Chem. Solids 1985, 46, 309−320. (11) Guo, X.; Waser, R. Space charge concept for acceptor-doped zirconia and ceria and experimental evidences. Solid State Ionics 2004, 173, 63−67. (12) Guo, X.; Zhang, Z. Grain size dependent grain boundary defect structure: Case of doped zirconia. Acta Mater. 2003, 51, 2539−2547. (13) Kjølseth, C.; Fjeld, H.; Prytz, O.; Dahl, P. I.; Estournès, C.; Haugsrud, R.; Norby, T. Space-charge theory applied to the grain boundary impedance of proton conducting BaZr0.9Y0.1O3−δ. Solid State Ionics 2010, 181, 268−275. (14) Göbel, M. C.; Gregori, G.; Maier, J. Size effects on the electrical conductivity of ceria: Achieving low space charge potentials in nanocrystalline thin films. J. Phys. Chem. C 2013, 117, 22560−22568. (15) Lin, Y.; Fang, S.; Su, D.; Brinkman, K. S.; Chen, F. Enhancing grain boundary ionic conductivity in mixed ionic−electronic conductors. Nat. Commun. 2015, 6, 68241−68249. (16) Guo, X.; Waser, R. Electrical properties of the grain boundaries of oxygen ion conductors: acceptor-doped zirconia and ceria. Prog. Mater. Sci. 2006, 51, 151−210. (17) Guo, X.; Sigle, W.; Maier, J. Blocking grain boundaries in Yttriadoped and undoped ceria ceramics of high purity. J. Am. Ceram. Soc. 2003, 86, 77−87. (18) Tschöpe, A.; Kilassonia, S.; Birringer, R. The grain boundary effect in heavily doped cerium oxide. Solid State Ionics 2004, 173, 57− 61. (19) Sato, K.; Shanai, D.; Hotani, Y.; Ougizawa, T.; Ito, K.; Hirata, K.; Kobayashi, Y. Positronium formed by recombination of positronelectron pairs in polymers. Phys. Rev. Lett. 2006, 96, 2283021− 2283024. (20) Sato, K.; Murakami, H.; Ito, K.; Hirata, K.; Kobayashi, Y. Probing the elemental environment around the free volume in polymers with positron annihilation age-momentum correlation spectroscopy. Macromolecules 2009, 42, 4853−4857. (21) Sato, K.; Baier, F.; Sprengel, W.; Würschum, R.; Schaefer, H.-E. Study of an order-disorder phase transition on an atomic scale: The example of decagonal Al-Ni-Co quasicrystals. Phys. Rev. Lett. 2004, 92, 1274031−1274034. (22) Sato, K.; Murakami, H.; Sprengel, W.; Schaefer, H.-E.; Kobayashi, Y. Nanocrystallization mechanism of amorphous Fe78B13Si9. Appl. Phys. Lett. 2009, 94, 1719041−1719043. (23) Sato, K. Grain-boundary structures associated with ionic transport in Gd-doped ceria nanostructured electrolyte. J. Phys. Chem. C 2015, 119, 5734−5738. (24) Shirpour, M.; Gregori, G.; Merkle, R.; Maier, J. On the proton conductivity in pure and gadolinium doped nanocrystalline cerium oxide. Phys. Chem. Chem. Phys. 2011, 13, 937−940. (25) Sato, K.; Fujimoto, K.; Kawamura, K.; Dai, W.; Hunger, M. Rheological mechanism of long-term self-assembly in saponite nanoparticles. J. Phys. Chem. C 2012, 116, 22954−22959. (26) Sato, K.; Fujimoto, K.; Dai, W.; Hunger, M. Molecular mechanism of heavily adhesive Cs: why radioactive Cs is not decontaminated from soil. J. Phys. Chem. C 2013, 117, 14075−14080. (27) Kirkegaard, P.; Eldrup, M. Positronfit extended: A new version of a program for analysing position lifetime spectra. Comput. Phys. Commun. 1974, 7, 401−409.

(28) Yoshida, H.; Deguchi, H.; Miura, K.; Horiuchi, M.; Inagaki, T. Investigation of the relationship between the ionicconductivity and the local structures of singly and doubly doped ceria compounds using EXAFS measurement. Solid State Ionics 2001, 140, 191−199. (29) Li, P.; Wei Chen, I.; Penner-Hahn, J. E. X-ray-absorption studies of zirconia polymorphs: I. Characteristic structures. Phys. Rev. B: Condens. Matter Mater. Phys. 1993, 48, 10063−10073. (30) Keeble, D. J.; Wicklein, S.; Dittmann, R.; Ravelli, L.; Mackie, R. A.; Egger, W. Identification of A- and B-site cation vacancy defects in perovskite oxide thin films. Phys. Rev. Lett. 2010, 105, 2261021− 2261024. (31) Č ížek, J.; Melikhova, O.; Procházka, I.; Kuriplach, J.; Kužel, R.; Brauer, G.; Anwand, W.; Konstantinova, T. E.; Danilenko, I. A. Defect studies of nanocrystalline zirconia powders and sintered ceramics. Phys. Rev. B: Condens. Matter Mater. Phys. 2010, 81, 0241161− 02411619. (32) Sato, K. Densification dynamics of Gadolinium-doped Ceria upon sintering. Jpn. J. Appl. Phys. 2012, 51, 0773011−0773014. (33) Sato, K.; Baier, F.; Rempel, A. A.; Sprengel, W.; Schaefer, H.-E. Observation of high-temperature thermal vacancies in Al70Pd21Mn9 quasicrystal. Phys. Rev. B: Condens. Matter Mater. Phys. 2003, 68, 2142031−2142034. (34) Sato, K.; Uchiyama, H.; Takahashi, Y.; Kanazawa, I.; Suzuki, R.; Ohdaira, T.; Takeuchi, T.; Mizuno, T.; Mizutani, U. Positronannihilation studies of rhombic triacontahedral-type icosahedral quasicrystals and their 1/1 and 2/1 approximants in Al-Mg-Zn alloy system. Phys. Rev. B: Condens. Matter Mater. Phys. 2001, 64, 0242021− 0242028. (35) Krause-Rehberg, R.; Leipner, H. S. Positron Annihilation in Semiconductors; Springer-Verlag: Berlin, 1999.

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DOI: 10.1021/acs.jpcc.7b07978 J. Phys. Chem. C XXXX, XXX, XXX−XXX