Atomic Scale Origin of Enhanced Ionic Conductivity at Crystal Defects

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Atomic-scale origin of enhanced ionic conductivity at crystal defects Bin Feng, Ryo Ishikawa, Akihito Kumamoto, Naoya Shibata, and Yuichi Ikuhara Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.9b00506 • Publication Date (Web): 20 Feb 2019 Downloaded from http://pubs.acs.org on February 21, 2019

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Atomic-scale origin of enhanced ionic conductivity at crystal defects Bin Feng†*, Ryo Ishikawa†, Akihito Kumamoto†, Naoya Shibata†,‡ and Yuichi Ikuhara†,‡

†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

ACS Paragon Plus Environment

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ABSTRACT: In oxide materials, the presence of dislocations often strongly affects the properties of thin film and multilayer devices. Although it was reported that ionic conduction can be enhanced by introducing dislocations in ionic conductors, the underlying mechanism of such enhancement remains unclear. Here we analyzed the ionic conduction enhancement in an yttria-stabilized zirconia (YSZ) single edge dislocation from structural point of view, using atomic resolution scanning transmission electron microscopy (STEM). First, atomic structure and chemistry of a dislocation in YSZ were characterized by STEM and energy dispersive X-ray spectroscopy (EDS). A relative ionic conduction variation map around the dislocation was then estimated based on the well-established strain-conductivity and chemistry-conductivity relationships in YSZ. We propose that a faster ionic conductivity path can be formed around the dislocation core due to the coupling of the tensile strain field and dopant segregation, which could account for enhanced ionic conductivity along dislocations.

KEYWORDS: dislocations, ionic conductivity, yttria-stabilized zirconia (YSZ), aberrationcorrected STEM, atomic-resolution EDS


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Solid oxide fuel cells (SOFCs) are one of the most promising technologies to directly convert chemical energies into electricity, with high efficiency and low emission of wasted pollutions1-4. For future worldwide commercialization, it is highly desirable to reduce the current working temperature (~ 1000 °C) down to the intermediated temperature range (~ 600 °C)5. Much efforts have been devoted to optimize conventional oxide electrolytes such as trivalent rare earth (RE) doped ZrO2 and CeO2 with a fluorite-type structure. To date, two types of modifications have been widely used for the optimization of ionic conductors6-10: doping trivalent RE atoms (oxygen vacancy doping) and artificially manipulating strain fields. These effects for ionic conductivity I can be simply interpretable from Arrhenius equation of 𝐼0

(

𝜀𝑎

)

𝐼 = 𝑇 exp ― 𝑘𝐵𝑇 ,

(1)

The carrier concentration tailored by doping oxygen vacancies mainly contributes to the preexponential factor I0; and the tensile strain field mainly contributes to the reduction of activation energy for oxygen migration (εa), since the expansion of lattice constant increases the oxygen migration space and reduce the bonding strength between oxygen and cations9. The optimization


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of the above relationship traditionally relies on the engineering of dopant species and their concentrations, resulting in an enhancement of ionic conductivity up to one order of magnitude8. On the other hand, recent works show that elastic strain field induced by the lattice size mismatch at the artificially introduced hetero-interfaces could improve conductions for more than two orders of magnitudes9. Such enhancement can be realized in thin films and multilayers9-16, which is recognized as a new design principle for high performance electrolyte materials. However, in thin films and multilayers, a large density of dislocations are usually formed at film-substrate interfaces17 and inside the films18, to accommodate the mismatch between film and substrate. Therefore, understanding the role of dislocations on transport properties is of paramount importance for the future design of high-performance electrolyte materials. It has been widely reported that dislocation core can be fast diffusion path in both metals19 and oxide materials20-21, known as pipe diffusion. In addition, dislocations possess surrounding longrange strain fields and resultant local chemistry22, which may also change the local ionic conductivity as shown in Eq. (1). Therefore, it is natural to consider that the ionic conductivity along dislocations might be largely altered compared with that of bulk. Several experimental studies have indeed reported faster oxygen ion diffusions along dislocations in electrochemical materials. Sillassen et al. tried to clarify the role of dislocations by using thin films13, throughout misfit dislocation network formed at these film-substrate interface. They found an enhanced ionic conductivity for yttria-stabilized zirconia (YSZ) epitaxial film on MgO, and they suggested that the enhancement is due to the misfit dislocations and interface misfit strain. Navickas et al. found that oxygen diffusion was accelerated along threading dislocations in LaSrMnO3 thin film23, of which they attributed the enhancement to the higher oxygen vacancy concentration in the dislocation induced by Sr segregation. Otsuka et al. built up a well-defined model system to


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analyze the dislocation effect: they introduced well-aligned, high-density dislocations in YSZ single crystal by uniaxial compression and measured the ionic conductivity. It was shown that the total ionic conductivity can be enhanced by the introduction of high-density of dislocations24, and the ionic conduction enhancement from each single dislocation was estimated based on the dislocation density change. They proposed pipe diffusion through dislocation cores accounts for the high ionic conduction along dislocation cores24. On the other hand, however, Otsuka et al. also found that such enhancement gradually disappears at higher temperatures when exceeding 850 °C (Ref 24). This agrees with the theoretical calculation study reported by Sun et al. for trivalent RE doped CeO2 dislocation at 900 °C (Ref 25). They suggested that the oxygen diffusion is even slightly slowed down in the dislocation core at this high temperature, due to the excessive segregation of RE dopants which strengthens the defect-defect interactions and reduces the oxygen diffusivity25. These results indicate that enhancement of dislocation conductions in those electrochemical materials is complicated, which cannot be simply explained by the conventional pipe diffusion behavior along the dislocation core. In summary, despite a number of experimental findings of fast oxygen diffusion along dislocations, the fundamental mechanism of such enhancement is still not fully clarified. A thorough understanding on the role of dislocations necessitates experimental information on the local ionic conduction of a single dislocation. However, it is still technically impossible to perform such measurement at single dislocation level due to a lack of spatial resolution. Even state-of-the-art scanning probe microscopy only allows local conductivity measurement at about 20 nm spatial resolution26, while the actual dislocation area is smaller by one order of magnitude. In this study, we attempt to investigate the ionic conductivity change at one single edge dislocation in YSZ from the structural point of view, by locally quantifying the magnitude of strain and Y-


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concentration using high-angle annular dark-field (HAADF) scanning transmission electron microscopy (STEM) combined with energy dispersive X-ray spectroscopy (EDS) at atomic resolution. On the basis of previously reported strain-conduction and Y-concentration-conduction relationships in YSZ6, 9, 27, we quantitatively converted the atomic-resolution STEM images into relative ionic conductivity variation maps, providing a high spatial resolution estimation of relative ionic conductivity change around the dislocation. We found that the tensile-strained area strongly accelerates the oxygen ionic conduction, while Y segregation around the dislocation core slightly worsens the conduction. As a result, a dislocation can provide faster ionic conduction path, of which the enhancement depends on the temperature. These findings successfully explained the previous inconsistencies of whether and how the ionic conductivity changes along dislocations24,25,36, and based on these results, we proposed an atomic-scale structural origin of the enhanced ionic conduction around dislocation. We first fabricated a pure edge dislocation array along the low-angle tilt grain boundary by bicrystal method28. Two well-oriented (110)/[110] YSZ single crystals (18 Y cation-mol% doped ZrO2, where we will use cation-mol% for Y in the following text) with a 2.5° mistilt angle from the (110) plane each, were annealed at 1600 °C for 15 h. This approach provides us an ideal system for the dislocation analysis by STEM28-29, and a well-defined dislocation array is confirmed from the low-magnification HAADF-STEM image in the Fig. 1a. Fig. 1b shows a high-magnification image of a dislocation core region. Since the intensity in HAADF-STEM image is strongly related to the atomic number (Z-contrast)30, we can only image the position of cation atomic columns intermixed with both Zr (Z = 40) and Y (Z = 39), shown as the bright spots. By drawing the Burgers circuit around the core, it shows that an extra half plane is on the upper region of dislocation with the Burgers vector of b = a/2[-110], where a corresponds to the lattice constant of YSZ. The


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dislocations prepared here are the same as those obtained via high-temperature compressing tests reported by Otsuka et al.24, and the core structure is also consistent with the previous highresolution TEM study24, 28. Since Z-contrast imaging is not suitable to identify the distribution of Y (Z=39) or Zr (Z=40), we performed STEM-EDS spectrum imaging with a probe size of about 1.2 Å. Fig. 2a and b show the atomically resolved STEM-EDS maps of Zr and Y, respectively. The Y intensity in the opposite side of the extra half plane is slightly higher than that in the extra half plane region. We also found that O concentration in these area follows the same tendency as those inside the bulk: lower O concentration in the Y segregated area and higher O concentration in the Y depleted area, as shown in Fig. S1. Such chemical redistribution area and concentration change around the dislocation is small, which is different from the case in RE doped CeO2 predicted by calculations25. We also note that there is no space charge layer in these dislocations probably because of the structural features of the dislocation: although the dislocation possesses an extra half plane, the overall structure is atomically stoichiometric. As a result, the charge neutrality could be locally compensated by Y and O vacancies in the vicinity of dislocation core (Supplementary Information). It is also known that phase transformation occurs in YSZ as a function of Y concentration, however, the Y concentration variation around the dislocation core is considerably small so that these areas should maintain the cubic phase. Fig. 3a and b show the projected 2D strain tensor diagonal components of εxx and εyy around the dislocation core, calculated by geometric phase analysis (GPA)31-32 using HAADF-STEM image. The strain of the bulk region was selected as a reference area where the strain components are below 0.4%. The εxx strain in the upper region of the dislocation with the extra half plane shows negative, indicating compressive strain field in the upper part, while the tensile strain field is confirmed in the lower region of the dislocation core. In a similar manner, εyy is shown in the Fig.


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3b and the distribution shows elongated cross pattern. The strain fields obtained here are consistent with those predicted from isotropic elastic theory of a typical edge dislocation22. The strain fields also give rise to the segregation of Y: since the atomic radius of Y is larger than Zr, Y segregation can be driven by minimization of elastic strain due to the size mismatch, forming the so-called “Cottrell atmosphere”33. Although we have only access to the experimental projected second order strain tensor, it is predicted that εzz is 0 in theory22, so only the effect of εxx and εyy will be taken into account in the following discussions. After obtaining the atomic-scale Y-concentration and strain fields around the dislocation core, we then calculated ‘relative ionic conductivity change’ along the dislocation core. This can be achievable based on the previously reported Y-concentration-conductivity and strain-conductivity relationships6, 9, 27, of which the details are described in the Methods section. Fig. 4a and b show the ionic conductivity profiles as a function of Y-concentration Is(c,T)s=0% and strain Ic(s,T)c=18 cation-mol%,

where c is Y-concentration, s is strain magnitude, T = 800K (as an example) and the

subscriptions of s, c indicate the constant. We fitted polynomial function (up to 4th order for Ic(s,T) and 6th order for Is(c,T)) to the data points reported6, 9, 27 to get each conductivity profiles. For the following calculation, we referred the values of the fitted curves in Fig. 4a and b. We also reasonably assume that the Y segregation amount and strain field distribution (denote as c(r) and s(r), where r is 2D coordinates) remains the same at these temperatures, based on the fact that temperature range discussed in this study (< 1000 °C) is relatively low for cation diffusion34, and thermal lattice expansion is also small35. The Y-concentration map of c(r) is given in STEM-EDS map, and we obtained a 2D ionic conductivity map of Is(c(r),T) normalized by the bulk Yconcentration (Fig. 4c, T = 800K). One can clearly see that the ionic conduction is lessened by the Y segregation at the lower region of the dislocation core. Since the Y-concentration of the bulk


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region here is already in the ‘ionic conductivity maximum’ range (~18 cation-mol%), further Y segregation should lead to a decrease of ionic conductivity in the tensile strained area compared with that in the bulk; even with a slight depletion of Y, the concentration in compressive strained area is in ‘ionic conductivity maximum range’, so that there is no significant reduction of the ionic conductivity. Similar segregation-induced conduction degradation was reported from theoretical calculations for RE doped CeO2 dislocations25. However, from another point of view, faster conduction along dislocation would be achievable if the initial Y-concentration in the bulk matrix is lower within the concentration range of cubic phase, so that Y-concentration can reach ‘ionic conductivity maximum’ range via segregation. In a similar manner, we also estimate the contribution of local strain magnitude for the normalized ionic conductivity Ic(s(r),T) from the strain magnitude map of s(r), where we used total strength magnitude of (εxx+ εyy) obtained from GPA analysis. According to the profile of Fig. 4b, the ionic conductivity is enhanced at about 8%12% tensile strain region, indicating that the ionic conductivity is increased at the lower region of the dislocation core (open-ring-shape in Fig. 4d) compared with that in the bulk. Although the result of 800K is shown here as an example, all the analysis was carried at different temperatures independently in the same way, and the details can be found in Table 1. Note that according to the continuum dislocation theory, the strain field extends into the bulk area even far from the dislocation core. However, we show that the contributions from these areas are small enough to be neglected (Supplementary Information). Now we consider the coupling effect of both Y segregation and strain for the ionic conductivity at the dislocation. As a first order approximation, we adopted that the parameters of Is(c(r),T) and Ic(s(r),T) are independent. In addition, the convoluted effects between dopant enrichment and strain field are taken into account by introducing another parameter of α(c,s). Based on the


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viewpoint of Arrhenius equation given in Eq. (1), where Is(c(r),T) corresponds to the preexponential factor I0, and Ic(s(r),T) corresponds to the exponential term in Eq. (1), the total ionic conductivity at this approximation can be expressed as following: 𝐼(𝒓, 𝑐, 𝑠, 𝑇) = 𝐼𝑠(𝑐(𝒓), 𝑇) ∙ 𝐼𝑐(𝑠(𝒓), 𝑇) ∙ 𝛼(𝒄,𝒔) ∙ 𝐼𝑏𝑢𝑙𝑘(𝑐,𝑇) = 𝐼𝑒𝑓𝑓(𝒓, 𝑇) ∙ 𝐼𝑏𝑢𝑙𝑘(𝑐,𝑇)

(2)

Where we have got the Y-concentration and strain magnitude at the position r, and Ibulk(c,T) is the bulk ionic conductivity with a certain Y-concentration at certain temperature. Since α(c,s) is still not clearly understood, we simply assume α(c,s)=1 in this study as an approximation, where we will show in the following discussions that such simple assumption still enables us to explain all the experimental results. Based on Eq. (2), we obtained relative ionic conductivity variation Ieff at the dislocation, and the map is given in the Fig. 5 at different temperatures. Fig. 5 indicates that the dislocation ionic conductivity is always enhanced compared with that in the bulk at these temperatures, which agrees with Otsuka’s report24. Fig. 5 also shows that the enhanced area is very small (about 1 nm2 area). Thus it suggests that the total ionic conductivity enhancement can be detectable only when the dislocation density is high, which is in good agreement with our previous macroscopic ionic conductivity measurements24. Combining with the Fig. 4a and b, we found that the effect of strain is much stronger than that of Y segregation, and hence the ionic conductivity at the dislocation core is mainly dominated by the strain field, forming a bat-wing-shape region as faster oxygen conduction paths. Further analysis quantifies the conductivity change in the dislocation compared with the bulk at each temperature, summarized in Table 1. Both the conductivity degradation due to Yconcentration and the conductivity enhancement due to strain are more prominent in the low temperature range. Although the Y segregation amount and the strain field magnitude do not


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change as the temperature rises, both the ionic conductivity degradation induced by the Y segregation and the ionic conductivity enhancement induced by the strain become weaker at higher temperature, as reported previously6,9. Since the conductivity change in the dislocation is mainly dominated by the strain as discussed, the resultant dislocation conductance enhancement gradually weakened as the temperature rises (Table 1). Such temperature dependence agrees well with our previous experimental results, that there is almost no conductivity difference between dislocation and the bulk at higher temperatures when exceeding 1000 °C (Ref 36). In this way, the dislocation conduction can be largely accelerated, forming a fast conductivity path along dislocation line in the low temperature region, e.g. the enhancement compared to the bulk can be 100 times at 400K. Note that we did not obtain a quantitative agreement with the previous experimental results24, probably because of simplified convoluted effects of α(c,s) in this study. Further study is necessary for more quantitative and precise discussions. Nevertheless, all these tendencies are in good agreement with the previous macroscopic ionic conductivity measurements along dislocations24, 36,

and thus our present result provides atomic-scale structural origin of ionic conduction change

in the dislocation. Our results indicate that the change of ionic conductivities in YSZ dislocation is governed by the tensile strain field and Y segregation near the dislocation core. Note that unlike the generally accepted pipe diffusion phenomenon along dislocations reported in oxide materials19-21, we do not think the dislocation core itself contributes much to the enhancement in the present cases. We will explain such difference with the case of pipe diffusion in Al2O3 (Ref 20 and 21) as an example. In Al2O3, the oxygen vacancy concentration is of p.p.m. level in the bulk, so that the bulk diffusion is very slow. The oxygen vacancy concentration is higher in the dislocation core due to a reduced defect formation energy, which results in faster diffusion along dislocations. However, in the case of YSZ or doped CeO2, the oxygen vacancy


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concentration level in the bulk is high, so that the bulk diffusion is fast. Since we found oxygen vacancy concentration change around the dislocation core is relatively small (see Supplementary Information), the dislocation core may contribute little to the faster diffusion in the case of YSZ. Our idea is supported by the theoretical25 and experimental evidences36 reported by previous studies (see Supporting Information for a detailed discussion). All these facts further suggest that it should be the tensile strain field and Y segregation near the dislocation core, rather than dislocation core itself, that account for the enhancement of ionic conductivities in YSZ. In this way, simultaneous modification of both pre-exponential factor I0 and the activation energy εa in Eq (1) can be naturally realized in dislocations. “Fastest conduction along dislocations” can be engineered by optimizing the segregation amount and strain conditions, providing fast ionic conductivity paths even inside conventional materials at the intermediate-temperature-range. Our conclusion should be applicable to other oxide ion conductors such as trivalent doped CeO2. It should also be valid for other defects such as misfit dislocations in hetero-interfaces and threading dislocations inside thin films. Therefore, the contribution from these defects should be taken into account when discussing the ionic conductivity change in those strained thin films, for example, the misfit dislocation may also accelerate the in-plane ionic conductivity in strained thin films. Unlike stand-alone nanoparticles or substrate-supported thin films, dislocations are buried inside bulk matrix, which may provide fast oxygen diffusion paths within the crystal. This is important for device engineering when considering the mechanical strength and the adaptability with cathode or anode materials. In addition, since the dislocations are ubiquitous and technologically controllable recently29, they would be of fundamental interests and potentials for future intermediate-temperature-range device-designing strategies in fuel cell technologies.


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In summary, we directly observe the atomic structure, strain field and corresponding element distributions in a single YSZ edge dislocation by using atomic-resolution STEM imaging and spectroscopy. By calculating the ionic conductivity change in these regions, we estimated the spatial distribution of the ionic conductivity enhancement around the dislocation core. Our results suggest the coupling of tensile strain field and component change around the dislocation core can contribute to the enhanced oxygen ionic conduction, and such enhancement is depending on the temperature with higher increment in the lower temperature range. These results help us to consistently explain the past studies on the dislocation conductions, and therefore, providing atomistic insights into enhanced dislocation ionic conductivity in the intermediate-temperaturerange. We believe our electron-microscopy-based approaches can also be applied to investigate other physical or chemical property variations in localized areas such as defects, interfaces and nanocrystalline materials at atomic resolution. Methods Low angle tilt grain boundary in YSZ bicrystal with tilt angle of θ=5° was used for the present study. To fabricate such bicrystal, two YSZ single crystals (18 Y cation-mol% doped ZrO2) were first cut precisely by 2.5° from the (110) plane and polished. Then these two polished crystals were bonded under 1600 degree for 15 h in air. The bicrystal specimen for STEM observations was prepared by mechanical polishing and ion milling. The atomic structure of the edge dislocation in YSZ was characterized by aberration-corrected STEM (ARM200CF, JEOL Ltd.) operated at 200kV. The strain field in the dislocation was quantified by GPA analysis32, using a lower magnification STEM image. The strain of the bulk region was selected as a reference area where the strain components are below 0.5%. Element


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distribution was characterized by STEM-EDS. The STEM-EDS in the present study are equipped with two ultrasensitive silicon-drift detectors, and the solid angle is about 1.7 sr. The probe size was 1.2 Å with a probe current of about 60 pA. The total acquisition time was about 20 minutes with a dwell time of 10 µs per pixel. The samples are robust under the present experimental conditions37-38. Y (using Kα of 14.958 keV and Kβ of 16.738 keV) and Zr (using Kα of 15.776 keV and Kβ of 17.668 keV) signals were extracted with the selected EDS energy for each element. Note that due to the electron channeling effect using the present experimental condition, it is difficult to precisely determine the oxygen content with atomistic structural information39. However, the quantification of relative content between Y and Zr is little affected by such artifact since they occupy the same atom sites. In addition, we indeed found that the oxygen concentration is qualitatively decreased in the Y segregated area as shown in the Fig. S1, showing that the oxygen vacancies prefer to segregate with Y to preserve local charge neutrality, similar to the tendency in the bulk YSZ. Therefore, we used Y concentration mapping for the discussions in this study. The function of Y-concentration-conductivity and strain-conductivity were fitted using previously reported data6, 9, 27. For the Y-concentration-conductivity function, data points at each temperature were first extrapolated using the data from Ref 6, assuming a constant activation energy. Since all the comparison and discussion were made on the relative enhancement in those strained area compared with the bulk part at each temperature, this approximation would not affect our conclusions. For the strain-conductivity function, the data points for tensile-strain-conductivity was selected from Ref 9. Although many experimental studies based on thin films or superlattices were discussed for the correlation of strain-conductivity, too many factors were involved in those studies, including the quality of the films and defects inside. Moreover, these factors varies case by case, which results in huge scatter of the enhancement10. Therefore, we chose the data obtained


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by theoretical calculations from Ref 9, which well reproduces almost all of the experimental ionic conductivity enhancement reported so far10. On the other hand, little attentions were paid to quantify the effect of compressive strain since it is known to decrease the ionic conductivity. To the authors’ knowledge, the data reported by Ref 27 is the only quantitative data available for YSZ, so that these data points were used here. In addition, only two data points are available under compressive strain 27, we therefore arbitrary set the value of dot circled data (Fig. 4b) to be -1.67, the same value as those under 6% compressive strain reported from Ref 27. Since we only focus on the area under tensile strain with faster diffusivity in this study, any negative diffusivity value under compressive strain does not affect our conclusions. The data points in previous reports were obtained either under biaxial or isotropic strain condition9, 27, thus the sum of (εxx + εyy) was calculated and used for x-axis of strain function Ic(s,T) here. Although the strain field is not uniaxial in the dislocation, the lattice constant indeed increases in these tensile strained area, which should accelerate the oxygen diffusion in a similar way to the case of biaxial strain. Since the fractional change in volume around dislocation is (εxx+ εyy+ εzz), shear component of εxy can be neglected here. In addition, previous theoretical studies show that the strain effect on oxygen diffusivity is little affected by Y concentration change

40,

therefore we used the same strain-

conductivity function shown in Fig 4b in both the dislocation and bulk part here. After obtaining the Is(c(r),T) and Ic(s(r),T) maps in Figure 4, and the total conductivity maps in Figure 5, we summed up the enhancement pixel by pixel in the bat-wing-shape area shown in Figure 5, and the value is listed in the Table. Is(c), Ic(s) and Ieff in Table 1 are calculated from Is(c(r),T), Ic(s(r),T) and total conductivity maps, respectively.


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Figure 1. HAADF-STEM image of the edge dislocation taken from [110] direction. (a) Low magnification image, in which the dislocations are indicated by the arrows. (b) High magnification image. The Burgers vector of the edge dislocation is b = a/2[-110]


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Figure 2. Chemistry of a single edge dislocation in YSZ: Net count elemental mapping obtained by STEM-EDS for (a) Zr K and (b) Y K line. (c) HAADF STEM image of the corresponding area.


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Figure 3. Strain field obtained by GPA analysis based on HAADF-STEM image. (a) εxx. (b) εyy. (c)HAADF STEM image of the corresponding area.


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Figure 4. Relative conductivity change originated from (a), (c) component change, and (b), (d) strain state. (a) Is(c, T)s=0%as a function of Y-concentration obtained by fitting experimental data. The case of 800 K is shown for example. (b) Ic(s, T)c=18 Y% as a function of strain state obtained by fitting calculated data at 800 K. (c) Relative conductivity change


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Is(c(r), T)s=0%calculated from the component change. (d) Relative conductivity change Ic(s(r), T)c=18 Y% calculated from the strain change.


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Figure 5. Relative ionic conductivity change at the dislocation with the corresponding color scale. (a) 400 K, (b) 600 K, (c) 800 K and (d) 1000 K. (e) HAADF STEM image of the corresponding area.


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Table 1. Relative ionic conductivity change in the bat-wing area shown in Figure 4. The effects of the Y-concentration (Is(c)), strain field (Ic(S)) and their coupling (Ieff) on the ionic conductivities are listed separately.

Temperature (K)

Relative conductivity of bat-wing area to the bulk value

Is(c)

Ic(s)

Ieff

400

4 x 10-2

3 x 103

1 x 102

600

1 x 10-1

2 x 102

3 x 10

800

3 x 10-1

6 x 10

2 x 10

1000

4 x 10-1

3 x 10

1 x 10


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ASSOCIATED CONTENT

Supporting Information.

The supporting information includes details of experimental STEM-EDS data for oxygen distribution in the dislocation, discussions on the estimation of conductivity enhancement in the extended tensile strained area, and discussions on the effects of the dislocation core on the conductivity enhancement. The supporting information is available free of charge on the ACS Publication website.

AUTHOR INFORMATION

Corresponding Author *Bin Feng, [email protected]

ACKNOWLEDGMENT


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This work was supported by Grants-in-Aid for Specially promoted Research (Grant No. JP17H06094), Grant-in-Aid for Scientific Research (A) (JP 17H01316) and Grant-in-Aid for Early-Career Scientists (JP18K13982) from the Japan Society for the Promotion of Science (JSPS). A part of this work is also supported by a Grant-in-Aid for Scientific Research on Innovative Areas "Nano Informatics" (Grant No. 25106003) from JSPS and "Nanotechnology Platform" (Project No. 12024046) of the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan.

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Atomic Scale Origin of Enhanced Ionic Conductivity at Crystal Defects

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