Direct Observation of Oxygen Vacancy Distribution across Yttria

Oct 13, 2017 - Crystalline interfaces in materials often govern the macroscopic functional properties owing to their complex structure and chemical in...
0 downloads 3 Views 7MB Size
Download PDF
www.acsnano.org

Direct Observation of Oxygen Vacancy Distribution across Yttria-Stabilized Zirconia Grain Boundaries Bin Feng,† Nathan R. Lugg,† Akihito Kumamoto,† Yuichi Ikuhara,†,‡ and Naoya Shibata*,†,‡ †

Institute of Engineering Innovation, The University of Tokyo, 2-11-16 Yayoi, Bunkyo-ku, Tokyo 113-8656, Japan Nanostructures Research Laboratory, Japan Fine Ceramics Center, 2-4-1 Mutsuno, Atsuta-ku, Nagoya, Aichi 456-8587, Japan

‡

S Supporting Information *

ABSTRACT: Crystalline interfaces in materials often govern the macroscopic functional properties owing to their complex structure and chemical inhomogeneity. For ionic crystals, however, such understanding has been precluded by the debatable local anion distribution across crystal interfaces. In this study, using yttria-stabilized zirconia as a model material, the oxygen vacancy distribution across individual grain boundaries was directly quantified by atomic-resolution scanning transmission electron microscopy with ultrahigh-sensitive energy-dispersive X-ray spectroscopy. Combined with dynamical scattering calculations, we unambiguously show that the relative oxygen concentrations increase in four high-angle grain boundaries, indicating that the oxygen vacancies are actually depleted near the grain boundary cores. These results experimentally evidence that the long-range electric interaction is the dominant factor to determine the local point defect distribution at ionic crystal interfaces. KEYWORDS: grain boundaries, yttria-stabilized zirconia (YSZ), STEM-EDS, oxygen vacancy, interface chemistry

C

GB core and depletion of VO (with positive excess charge) in the adjacent SC layers.7 This model was indirectly supported by electrochemical impedance spectroscopy studies:7 the GB ionic conductivities are much lower than that of the bulk, which can be explained by the depletion of the mobile carrier of VO in the vicinity of the GBs. On the other hand, some theoretical studies have shown that VO segregates with Y into GBs without forming any VO depletion layers.12,15,16 In this scenario, the GB oxygen diffusion rate is supposed to be degraded due to the interface structural incoherency.17 Recently, direct experimental characterization of local structure and chemistry of YSZ GBs is becoming possible by scanning transmission electron microscopy (STEM) combined with chemical analyses such as energydispersive X-ray spectroscopy (EDS) and electron energy loss spectroscopy (EELS). Using these direct experimental techniques, Y segregation at GBs has been well confirmed.18−22 However, the experimental determination of O distribution across the GBs is still extremely challenging. Although a few reports claimed the segregation of VO at the GB cores, none of them show that oxygen concentrations increase in the adjacent SC layers, and thus the detailed O (VO) distributions across GBs are still under controversy both experimentally and theoretically.

rystalline interfaces such as grain boundaries (GBs), ubiquitous two-dimensional crystal defects inside materials, have a huge impact on the macroscopic properties of many technologically important materials.1−8 Owning to their local chemical inhomogeneity, GBs have exhibited unusual mechanical,1,2 electrical,3,4 and chemical properties5,8 which may not emerge in the bulk crystals. Detailed knowledge on how chemical inhomogeneity is induced at GBs should be essential to understand and control GB properties for future material design. The physical picture of interface chemistry in ionic crystals was established in detail based on the results of acceptor-doped ZrO2 and SrTiO3 GBs.7,9,10 Although this concept has been widely applied to a wide range of materials since then,9 interface chemistry is still controversial because several complex factors have been proposed to account for the final chemical distribution of GBs, such as atomic arrangements of GB cores,1,11 local point defect−defect interactions,12 and long-range electric interactions between point defects and GB cores.7 For instance, even in a simple ternary system of yttriastabilized zirconia (YSZ), a typical ionic material that is widely used as electrolyte material in solid oxide fuel cells (SOFCs),7,13 the applicability of thermodynamic models for interface defect chemistry14 is still under debate. The space charge (SC) theory predicts that the YSZ GB core is positively charged due to the intrinsic O vacancies (VO) and/or Zr interstitials. Such charged GBs lead to the segregation of Y (Zr4+ substituted by Y3+ equals negative excess charge) to the © 2017 American Chemical Society

Received: August 21, 2017 Accepted: October 13, 2017 Published: October 13, 2017 11376

DOI: 10.1021/acsnano.7b05943 ACS Nano 2017, 11, 11376−11382

Article

Cite This: ACS Nano 2017, 11, 11376-11382

Article

ACS Nano

Figure 1. HAADF-STEM images of the model GBs. (a) Σ5 (310), (b) Σ5 (210), (c) Σ9, and (d) Σ13 GB. Bright spots correspond to the cation (Zr and Y) atomic columns. The GB is well bonded at atomic scale. Scale bar, 1 nm.

Figure 2. EDS maps (net count) for the model GBs and corresponding intensity profiles. (a−d) Element map for Zr K in (a) Σ5 (310), (b) Σ5 (210), (c) Σ9, and (d) Σ13, with the corresponding average intensity line profiles (au). (e−h) Element map for Y K in (e) Σ5 (310), (f) Σ5 (210), (g) Σ9, and (h) Σ13, with the corresponding average intensity line profiles (au). (i−l) Element map for O K in (i) Σ5 (310), (j) Σ5 (210), (k) Σ9, and (l) Σ13, with the corresponding average intensity line profiles (au). The intensity profile (net counts) is normalized by the total counts, obtained by summing the X-ray counts in the direction parallel to the GB. Error bars indicate standard deviation. The GBs are indicated by the arrows. Scale bar for EDS maps, 10 nm.

We refer, respectively, to these GBs as Σ5(210), Σ5(310), Σ9, and Σ13 GBs in the following text. By systematic analysis of GB geometry and the resultant electron channeling effect in STEMEDS, quantitative oxygen chemical composition maps in the vicinity of the GBs were obtained. These results should answer the long-standing question on how VO distributes across the ionic crystalline interfaces.

In this study, we report the direct observation of local elemental distributions in the vicinity of YSZ GBs by using aberration-corrected STEM-EDS. Four coincidence site lattice (CSL) model GBs were fabricated by a bicrystal method. Here, Σ5[001]/(210), Σ5[001]/(310), Σ9[110]/(221), and Σ13[001]/(510) GBs, where Σ denotes the degree of geometrical coincidence of crystalline interfaces,15 are selected. 11377

DOI: 10.1021/acsnano.7b05943 ACS Nano 2017, 11, 11376−11382

Article

ACS Nano

Figure 3. EDS concentration maps (atomic %) for the model GBs and corresponding intensity profiles. (a−d) Element map for Zr K in (a) Σ5 (310), (b) Σ5 (210), (c) Σ9, and (d) Σ13, with the corresponding relative concentration change. (e−h) Element map for Y K in (e) Σ5 (310), (f) Σ5 (210), (g) Σ9, and (h) Σ13, with the corresponding relative concentration change. (i−l) Element map for O K in (i) Σ5 (310), (j) Σ5 (210), (k) Σ9, and (l) Σ13, with the corresponding relative concentration change. The relative component change (atomic %) is obtained by averaging the component in the direction parallel to the GB. Error bars indicate standard deviation. The GBs are indicated by the arrows. Scale bar for EDS maps, 10 nm.

RESULTS AND DISCUSSION Figure 1 shows high-angle annular dark-field (HAADF) images of the GBs fabricated in this study. All the GBs were well bonded at the atomic level, and the atomic structures of these GBs are consistent with previous reports.19,20,23,24 The local GB chemistry was then investigated by STEM-EDS. To reduce the effects of electron channeling and gain a proper insight into the GB chemistry (which will be discussed later), all these GBs were set to be about 15° tilted from the zone axis orientation while maintaining interface edge-on conditions. Figure 2 shows the STEM-EDS elemental maps for the Zr K lines, Y K lines, and O K lines in different GBs, respectively. The GB positions are indicated by the arrows. To access any small local chemistry changes at the GBs, we formed line profiles across the GBs from the maps by integrating the signal parallel to the GBs. The line profiles are also shown in Figure 2. The total counts for Zr and O decreased in the vicinity of the GB (compared with the bulk), while the Y counts slightly increased in the Σ9 GB but remained almost unchanged in the other GBs. At first glance, one might conclude that Y is increased and O is decreased in the vicinity of the GBs from these results, consistent with the previous reports that both Y and VO segregate to GBs.19,20,24 However, the quantity we are interested in is how the relative amounts of Y and O change across the GBs, which cannot be directly estimated from the EDS counts maps. Increase (or decrease) in EDS counts at the GBs cannot be used to estimate the relative amounts of Y and O, because there is a possibility of

density and/or thickness changes at the core of the GBs. For instance, the previous static lattice calculations have shown the effect of GB geometry on the GB core densities: the density of both O and Zr is lower near the GBs even in a stoichiometric ZrO2 GB model without any O nor Zr vacancies.25 In this case, the lower signal in the GB merely originates from the relatively nondense atomic configurations (compared with the bulk). In addition, preferential ion milling may locally decrease the total specimen thickness at the GB cores, which might also decrease the EDS counts from the GB regions. Therefore, the relative changes in concentration at GBs cannot be simply deduced from the raw EDS count maps. Here, we interpret the GB chemistry from their component change, defined by the atomic ratio, namely, Ni/Ntotal, where Ni is the number of atoms for element i and Ntotal is the total number of atoms. The atomic ratio of each element was calculated by the Cliff−Lorimer method,26,27 and the resultant maps are shown in Figure 3. The corresponding line profiles across the GBs are also shown. It is now clear that Y segregates (and correspondingly Zr depletes) at all the high-angle GBs studied here. Moreover, the Y segregation amount is dependent on the GB character, which is consistent with the previous study.19 On the other hand, while the concentration of Y (NY/ (NZr + NY + NO)) increased and Zr (NZr/(NZr + NY + NO)) decreased across all the GBs, the O (NO/(NZr + NY + NO)) shows higher concentration at the GB cores than in the bulk. These results indicate that O concentration is increased at the 11378

DOI: 10.1021/acsnano.7b05943 ACS Nano 2017, 11, 11376−11382

Article

ACS Nano GB cores. In other words, VO is depleted around the GBs. Such spatial distribution of O concentration is found in all the GBs studied here, suggesting that the increase in O concentration in the vicinity of GBs should be a general feature of YSZ highangle GBs. To confirm these compositional changes do originate from the GB structures rather than from the artifacts of complex electron scattering effects, we have quantitatively considered the effect of electron channeling using multislice STEM-EDS simulations.28 The Σ9 GB with the highest compositional variation among all the GBs studied here was selected for the model simulation, whose GB atomic structure has been previously well characterized.19 The model GB cell consists of only Zr and O with the stoichiometric composition. Such a model system should enable us to exclude the influence of any component changes on EDS signals, and the effect of channeling caused by the GB geometry can be properly evaluated. Simulations of both the on-axis model and 15° tilted edge-on model were systematically performed. Here, we only show the tilted results, which are consistent with the present experimental setup. Figure 4a and c show the simulated Zr K line and O K line STEM-EDS maps across the tilted GB. It can be seen that both Zr K (Figure 4b) and O K counts (Figure 4d) drop at the core of the GB. However, the calculated concentration profile for Zr (NZr/(NZr +NO)) and O (NO/ (NZr +NO)) shows different behavior as shown in Figure 4e: the Zr concentration slightly increases at the GB core, while O

slightly decreases at the same place. These results suggest that the electron channeling effect under the present experimental conditions may slightly underestimate the oxygen concentration at the core of the GB, although the stoichiometry is maintained in the present simulation cell. In the actual experiments, we clearly observed the increase in oxygen concentration at the core of the GBs. Combined with these simulations, we conclude that the increase of oxygen concentration found in the experiment must reflect the intrinsic GB structures. It is worth mentioning that, under the zone axis condition, which is commonly used for interface analysis, our simulation results show that the O concentration would be overestimated due to the complex electron scattering effect in GB (Supporting Information). The oxygen concentration indeed experimentally increased under zone axis condition; however, under such a condition, it is currently impossible to discriminate whether the increased O concentration is originated from the intrinsic GB structures or the artifact of electron channeling effects (Supplementary Discussion 1). These results call attention to the fact that the EDS count maps are not necessarily coincident with the actual concentration changes at GB cores under strong electron channeling conditions. In the bulk YSZ, VO has been considered to be stabilized near Y3+ ions (second-nearest neighbor) to maintain charge neutrality and to relax the strain.29 However, from our analysis shown above, Y concentration is increased while VO is decreased in the vicinity of GB cores, indicating that the elemental distribution near the GBs does not simply follow this mechanism. Here, we discuss the origin of such chemical distributions across the GBs. It is known that the atomic bonding environment in the GB cores significantly differs from those inside perfect crystals: forming coordination deficient sites, dangling bonds, and nonstoichiometric terminations.1,3 These structural factors indeed attract Y3+ to segregate to the GB.15,16,19,21 If such structural factors were the only reasons for the Y segregation, VO should also segregate to GB cores because VO formation can be also promoted by Y segregation and structural disorders.21,29 A schematic concentration profile of this situation is shown in Figure 5a. Our experimental results, however, do not agree with this profile. There should be other mechanisms to repel VO’s away from the GB cores. The plausible mechanism is the long-range Coulomb interaction between GB cores and VO’s, provided that the GB core is positively charged. Such a mechanism has been proposed through the classical SC theory. In this theory, an electrostatic potential difference between GB cores and bulk due to the enrichment of intrinsic immobile VO’s at the cores7 can electrically affect the point defect distributions across the GBs: negatively charged Y3+ ions are enriched in the GB core, while the mobile positively charged VO’s are depleted in the areas adjacent to the GB core.7 Such an SC model is schematically shown in Figure 5b,7 which reasonably explains the observed Y enrichment and VO depletion in the vicinity of GBs. Thus, we conclude that long-range electric interaction should be the dominant factor to determine the local point defect distribution across the YSZ GBs. However, in contrast to the SC model,7 the experimental profiles (i.e., the Σ9 GB) in Figure 5d show that both Y and O are increased even at the core of GBs. In this case, there should be no intrinsic immobile VO’s at the GB cores that should be responsible for the GB core charges. Thus, the origin of the positive core charge should be reconsidered. Here, we propose that the origin of GB extra charge might

Figure 4. Simulated EDS maps and corresponding composition profiles for the Σ9 GB. (a, c) Simulated EDS count map for (a) Zr and (c) O. (b, d) Corresponding integrated line profiles across the GB from the EDS map for (b) Zr K and (d) O K. (e) Element concentration calculated from the simulation. 11379

DOI: 10.1021/acsnano.7b05943 ACS Nano 2017, 11, 11376−11382

Article

ACS Nano

Figure 5. Schematic models for element distribution near GBs. (a) The GB is not charged in this mode, and Y segregation occurred to relax the strain. The VO’s segregate with Y, to maintain the local charge neutrality as those inside the bulk YSZ. Therefore, the Y concentration increases and O concentration decreases in GBs (and thus VO concentration increases). (b) SC model. The GB is assumed to be charged due to the high concentration of intrinsic VO’s. Thus, the VO concentration is extremely high in the GB core part. The GB charges attract Y and repel VO’s. As a result, both the Y and O concentration increase in adjacent to GBs (and thus VO concentration increases). (c) Model proposed in this study. The GB is assumed to be charged due to the nonperiodic atomic arrangement. The GB charges attract Y and repel VO’s. In contrast with model b, both the Y and O concentration always increases in the core part and adjacent to GBs. (d) Experimental element concentration, with the example of the Σ9 GB.

effects should provide an opportunity to experimentally verify the formation of SC layers in such heterointerface systems.

originate from the GB core structure themselves, such as the incoherent atomic configurations in the GB cores. As shown in Figure 1, the GBs studied here showed very different atomic configurations compared with that of bulk crystal. It has been reported that nonstoichiometric terminations and/or the breaking cation−anion bonding distributions could lead to a local spatial charge inhomogeneity in both dislocations and GBs.11,30,31 Thus, in the present YSZ GBs, such structural factors could cause charge inhomogeneities at their cores. Such assumptions can be further supported from our previously studies of the YSZ Σ3 GB,21 as shown in Figure S3a−f. It shows that, unlike the GBs reported in this study, the Y segregation is accompanied by the VO segregation in the case of Σ3 GB. This GB has an atomically coherent GB configuration (Figure S3g) without nonstoichiometric terminations, and thus there should be very small extra charges in the GB core. In this case, charge neutrality can be maintained simply by Y and VO as in the bulk crystal. In contrast to the Σ3 GB, all the other GBs studied here possess much larger structural distortions, as reported previously.8,19 For example, in the case of the Σ9 GB, it has been shown that cation sites with lower coordination numbers are densely formed along the GB core.19 In this case, the GB may structurally possess charge-imbalanced regions, and thus the total charge neutrality should be maintained by their adjacent long-range SC layers. The schematic profile of this scenario is shown in Figure 5c, which is most consistent with our experimental observations. The models we proposed here inevitably suggest that GB extra charges could be strongly related to the GB characters because GB structures are strongly dependent on the GB characters. The equilibrium GB defect distributions should be thus strongly dependent on the GB characters (Supplementary Discussion 2). These tendencies are consistent with our experimental observations. However, to obtain further quantitative correlation between GB core structures and GB extra charges, a detailed first-principles calculation for the individual GBs is necessary, which would be a future research direction. Finally, the present findings may shed light on the true understanding of other oxide heterointerface phenomena reported previously. The origin of enhanced ionic conductivities at these oxide heterointerfaces has been conventionally explained by the space charge effects.32−34 However, there has been no direct evidence of the formation of SC layers at these oxide heterointerface so far. Our method of using quantitative STEM-EDS with careful consideration of electron channeling

CONCLUSIONS In summary, we used atomic-resolution STEM-EDS to directly quantify the local oxygen vacancy distribution across YSZ GBs at subnanometer dimensions. We directly observed that the oxygen vacancies are depleted near the GB cores. These results experimentally evidence that long-range electric interaction is the dominant factor to determine the local point defect distribution across YSZ GBs. Our combined experimental and theoretical approaches using STEM-EDS should be a powerful way for directly investigating local interface chemistry in many oxide materials and devices. METHODS Bicrystal Sample Preparation. Four coincidence site lattice GBs were fabricated using diffusion bonding of two well-defined single crystals, resulting in a bicrystal of Σ5[001]/(210), Σ5[001]/(310), Σ9[110]/(221), and Σ13[001]/(510) GBs. The YSZ bicrystals containing model GBs were first fabricated by joining two YSZ single crystals at 1600 °C for 15 h in air.19 TEM specimens were prepared by mechanical polishing and Ar ion-beam milling. The thicknesses of the samples observed were about 30 nm. STEM-EDS Characterization and Simulations. GBs were observed using an STEM (JEM-ARM200CF, JEOL Co. Ltd.) operated at 200 keV. STEM-EDS mappings were acquired by scanning the beam over a wide field of view containing a GB, using the NORAN System 7 spectral analyzer and NSS3 spectral analysis software developed by Thermo Fisher Scientific Inc. In the present STEM-EDS characterization, we used a fine electron probe; thus it should be emphasized that the quantitative understanding of these results is difficult, due to the complexity of the electron channeling effect inside the crystal.35 To reduce the effect of electron channeling, we tilted the sample36 by 15 degrees while maintaining edge-on conditions for GBs. The EDS simulations were then performed to distinguish the effect of electron channeling from the effect of the intrinsic GB chemistry. The EDS collection solid angle was about 0.8 sr. The probe size was 1.2 Å with a probe current of about 60 pA. The samples are robust under the present experimental conditions.21 Using a top-hat filter for background subtraction, net counts for elemental maps were extracted with selected EDS energy for each element: O (using Kα of 0.525 keV), Zr (using Kα of 15.776 keV and Kβ of 17.668 keV), and Y (using Kα of 14.958 keV and Kβ of 16.738 keV) were summed, and the combined map is shown in Figure 2 and Figure 3. We experimentally obtained reference X-ray spectra for the background subtraction from the thin area of ZrO2, Y2O3, and Al2O3 powder for Zr, Y, and O, respectively. Theoretical k-factors with consideration for ionization cross sections 11380

DOI: 10.1021/acsnano.7b05943 ACS Nano 2017, 11, 11376−11382

Article

ACS Nano and detector geometry were used,27 and the composition calculated in the bulk part is consistent with the single-crystal YSZ used in this study. In order to acquire sufficient signal for the Zr and Y K lines, the total acquisition time was about 4 h with a dwell time of 10 μs per pixel. To examine the complex electron scattering effects in the GBs, we have performed multislice STEM-EDS simulations. Thermal diffuse scattering was included using an absorptive model. The crystal tilt was included via appropriate modifications to the multislice transmission function and propagator.36,37 An aberration-free probe with a semiconvergence angle of 25 mrad and accelerating voltage of 200 kV was assumed. Since these parameters correspond to an atomically sharp electron probe, the image simulations were first performed at atomic resolution (with a pixel spacing about 0.25 Å) and then binned to the larger pixel size (about 1.7 Å). This is consistent with the way the Thermo Fisher Scientific control system (we have used in the present study) continuously scans the beam across the specimen to obtain EDS signals. A Gaussian blur of 0.5 Å was applied to the final image to take into account the finite source size of the probe.38,39 The image was scaled such that there was an average of 600 counts per pixel, and Poisson noise was then applied to each pixel. The specimen was assumed to be 30 nm thick.

Current Transport at Grain Boundaries in High-Tc Superconductors. Nature 2005, 435, 475−478. (5) Lee, W.; Jung, H. J.; Lee, M. H.; Kim, Y.-B.; Park, J. S.; Sinclair, R.; Prinz, F. B. Oxygen Surface Exchange at Grain Boundaries of Oxide Ion Conductors. Adv. Funct. Mater. 2012, 22, 965−971. (6) Shibata, N.; Findlay, S. D.; Azuma, S.; Mizoguchi, T.; Yamamoto, T.; Ikuhara, Y. Atomic-Scale Imaging of Individual Dopant Atoms in a Buried Interface. Nat. Mater. 2009, 8, 654−8. (7) 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. (8) Feng, B.; Sugiyama, I.; Hojo, H.; Ohta, H.; Shibata, N.; Ikuhara, Y. Atomic Structures and Oxygen Dynamics of CeO2 Grain Boundaries. Sci. Rep. 2016, 6, 20288. (9) Gregori, G.; Merkle, R.; Maier, J. Ion Conduction and Redistribution at Grain Boundaries in Oxide Systems. Prog. Mater. Sci. 2017, 89, 252−305. (10) Guo, X.; Fleig, J.; Maier, J. Determination of Electronic and Ionic Partial Conductivities of a Grain Boundary: Method and Applicaion to Acceptor-Doped SrTiO3. Solid State Ionics 2002, 154− 155, 564−569. (11) Graser, S.; Hirschfeld, P. J.; Kopp, T.; Gutser, R.; Andersen, B. M.; Mannhart, J. How Grain Boundaries Limit Supercurrents in HighTemperature Superconductors. Nat. Phys. 2010, 6, 609−614. (12) Lee, H. B.; Prinz, F. B.; Cai, W. Atomistic Simulations of Grain Boundary Segregation in Nanocrystalline Yttria-Stabilized Zirconia and Gadolinia-Doped Ceria Solid Oxide Electrolytes. Acta Mater. 2013, 61, 3872−3887. (13) Steele, B. C. H.; Heinzel, A. Materials for Fuel-cell Technologies. Nature 2001, 414, 345−352. (14) Kliewer, K.; Koehler, J. Space Charge in Ionic Crystals. I. General Approach with Application to NaCl. Phys. Rev. 1965, 140, A1226−A1240. (15) Yokoi, T.; Yoshiya, M.; Yasuda, H. Nonrandom Point Defect Configurations and Driving Force Transitions for Grain Boundary Segregation in Trivalent Cation Doped ZrO2. Langmuir 2014, 30, 14179−88. (16) Yoshiya, M.; Oyama, T. Impurity and Vacancy Segregation at Symmetric Tilt Grain Boundaries in Y2O3-Doped ZrO2. J. Mater. Sci. 2011, 46, 4176−4190. (17) Fisher, C. A.; Matsubara, H. Oxide Ion Diffusion Along Grain Boundaries in Zirconia: A Molecular Dynamics Study. Solid State Ionics 1998, 113−115, 311−318. (18) Ikuhara, Y.; Thavorniti, P.; Sakuma, T. Solute Segregation at Grain Boundaries in Superplastic SiO2-Doped TZP. Acta Mater. 1997, 45, 5275−5284. (19) Shibata, N.; Oba, F.; Yamamoto, T.; Ikuhara, Y. Structure, Energy and Solute Segregation Behaviour of [110] Symmetric Tilt Grain Boundaries in Yttria-stabilized Cubic Zirconia. Philos. Mag. 2004, 84, 2381−2415. (20) Lei, Y.; Ito, Y.; Browning, N. D.; Mazanec, T. J. Segregation Effects at Grain Boundaries in Fluorite-Structured Ceramics. J. Am. Ceram. Soc. 2002, 85, 2359−2363. (21) Feng, B.; Yokoi, T.; Kumamoto, A.; Yoshiya, M.; Ikuhara, Y.; Shibata, N. Atomically Ordered Solute Segregation Behaviour in An Oxide Grain Boundary. Nat. Commun. 2016, 7, 11079. (22) Frechero, M. A.; Rocci, M.; Sanchez-Santolino, G.; Kumar, A.; Salafranca, J.; Schmidt, R.; Diaz-Guillen, M. R.; Dura, O. J.; RiveraCalzada, A.; Mishra, R.; Jesse, S.; Pantelides, S. T.; Kalinin, S. V.; Varela, M.; Pennycook, S. J.; Santamaria, J.; Leon, C. Paving The Way to Nanoionics: Atomic Origin of Barriers for Ionic Transport Through Interfaces. Sci. Rep. 2015, 5, 17229. (23) Dickey, E. C.; Fan, X.; Pennycook, S. J. Structure and Chemistry of Yttria-stabilized Cubic-zirconia Symmetric Tilt Grain Boundaries. J. Am. Ceram. Soc. 2001, 84, 1361−1368. (24) An, J.; Park, J. S.; Koh, A. L.; Lee, H. B.; Jung, H. J.; Schoonman, J.; Sinclair, R.; Gur, T. M.; Prinz, F. B. Atomic Scale Verification of Oxide-Ion Vacancy Distribution Near A Single Grain Boundary in YSZ. Sci. Rep. 2013, 3, 2680.

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.7b05943. Details of experimental and simulations of STEM-EDS for Σ9 GB with on-zone axis condition, experimental STEM-EDS for the Σ3 GB, and discussion of GB dependence on interface chemistry (PDF)

AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. ORCID

Bin Feng: 0000-0002-4306-2979 Naoya Shibata: 0000-0003-3548-5952 Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work is supported by a Grant-in-Aid for Scientific Research on Innovative Areas “Nano Informatics” (Grant No. 25106003) from Japan Society for the Promotion of Science (JSPS) and “Nanotechnology Platform” (Project No. 12024046) of the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan. This work was partially supported by Grants-in-Aid for Specially Promoted Research (Grant No. JP17H06094) and a Grant-in-Aid for Scientific Research (A) (JP 17H01316) from the JSPS. REFERENCES (1) Buban, J. P.; Matsunaga, K.; Chen, J.; Shibata, N.; Ching, W. Y.; Yamamoto, T.; Ikuhara, Y. Grain Boundary Strengthening in Alumina by Rare Earth Impurities. Science 2006, 311, 212−215. (2) Nie, J.; Zhu, Y. M.; Liu, J. Z.; Fang, X. Y. Periodic Segregation of Solute Atoms in Fully Coherent Twin Boundaries. Science 2013, 340, 957−960. (3) Sato, Y.; Buban, J. P.; Mizoguchi, T.; Shibata, N.; Yodogawa, M.; Yamamoto, T.; Ikuhara, Y. Role of Pr Segregation in Acceptor-State Formation at ZnO Grain Boundaries. Phys. Rev. Lett. 2006, 97, 106802. (4) Klie, R. F.; Buban, J. P.; Varela, M.; Franceschetti, A.; Jooss, C.; Zhu, Y.; Browning, N. D.; Pantelides, S. T.; Pennycook, S. J. Enhanced 11381

DOI: 10.1021/acsnano.7b05943 ACS Nano 2017, 11, 11376−11382

Article

ACS Nano (25) Oyama, T.; Yoshiya, M.; Matsubara, H.; Matsunaga, K. Numerical Analysis of Solute Segregation at Σ5 (310)/[001] Symmetric Tilt Grain Boundaries in Y2O3-Doped ZrO2. Phys. Rev. B: Condens. Matter Mater. Phys. 2005, 71, 224105. (26) Cliff, G.; Lorimer, G. W. The Quantitative Analysis of Thin Specimens. J. Microsc. 1975, 103, 203−207. (27) Lorimer, G. W. Quantitative X-Ray Microanalysis of Thin Specimens in The Transmission Electron Microscope: A Review. Mineral. Mag. 1987, 51, 49−60. (28) Allen, L. J.; D’Alfonso, A. J.; Findlay, S. D. Modelling The Inelastic Scattering of Fast Electrons. Ultramicroscopy 2015, 151, 11− 22. (29) Zacate, M. O.; Minervini, L.; Bradfield, D. J.; Grimes, R. W.; Sickafus, K. E. Defect Cluster Formation in M2O3 -Doped Cubic ZrO2. Solid State Ionics 2000, 128, 243−254. (30) Shibata, N.; Chisholm, M. F.; Nakamura, A.; Pennycook, S.; Yamamoto, T.; Ikuhara, Y. Nonstoichiometric Dislocation Cores in αAlumina. Science 2007, 316, 82−85. (31) Tochigi, E.; Kezuka, Y.; Nakamura, A.; Nakamura, A.; Shibata, N.; Ikuhara, Y. Direct Observation of Impurity Segregation at Dislocation Cores in An Ionic Crystal. Nano Lett. 2017, 17, 2908. (32) Sata, N.; Eberman, K.; Eberl, K.; Maier, J. Mesoscopic Fast Ion Conduction in Nanometer-scale Planar Heterostructures. Nature 2000, 408, 946−949. (33) Garcia-Barriocanal, J.; Rivera-Calzada, A.; Varela, M.; Sefrioui, Z.; Iborra, E.; Leon, C.; Pennycook, S.; Santamaria, J. Colossal Ionic Conductivity at Interfaces of Epitaxial ZrO2: Y2O3/SrTiO3 Heterostructures. Science 2008, 321, 676−680. (34) Yildiz, B. Stretching” The Energy Landscape of OxidesEffects on Electrocatalysis and Diffusion. MRS Bull. 2014, 39, 147−156. (35) Forbes, B. D.; D’Alfonso, A. J.; Williams, R. E. A.; Srinivasan, R.; Fraser, H. L.; McComb, D. W.; Freitag, B.; Klenov, D. O.; Allen, L. J. Contribution of Thermally Scattered Electrons to Atomic Resolution Elemental Maps. Phys. Rev. B: Condens. Matter Mater. Phys. 2012, 86, 024108. (36) Lugg, N. R.; Kothleitner, G.; Shibata, N.; Ikuhara, Y. On The Quantitativeness of EDS STEM. Ultramicroscopy 2015, 151, 150−159. (37) Chen, J. H.; VanDyck, D.; OpdeBeeck, M. Multislice Method for Large Beam Tilt with Application to HOLZ Effects in Triclinic and Monoclinic Crystals. Acta Crystallogr., Sect. A: Found. Crystallogr. 1997, 53, 576−589. (38) LeBeau, J.; Findlay, S.; Allen, L.; Stemmer, S. Quantitative Atomic Resolution Scanning Transmission Electron Microscopy. Phys. Rev. Lett. 2008, 100, 206101. (39) Dwyer, C.; Erni, R.; Etheridge, J. Measurement of Effective Source Distribution and Its Importance for Quantitative Interpretation of STEM Images. Ultramicroscopy 2010, 110, 952−957.

11382

DOI: 10.1021/acsnano.7b05943 ACS Nano 2017, 11, 11376−11382