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Identification of Actual Strain-induced Effect on Fast Ion Conduction in Thin Film Electrolyte Junsung Ahn, Ho Won Jang, Ho-Il Ji, Hyoungchul Kim, Kyung Joong Yoon, Ji-Won Son, Byung-Kook Kim, Hae-Weon Lee, and Jong-Ho Lee Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.7b04952 • Publication Date (Web): 09 Apr 2018 Downloaded from http://pubs.acs.org on April 9, 2018

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Identification of Actual Strain-induced Effect on Fast Ion Conduction in Thin Film Electrolyte Junsung Ahn,†, ‡ Ho Won Jang, ‡ Hoil Ji, † Hyoungchul Kim,† Kyung Joong Yoon,† Ji-Won Son†, § Byung-Kook Kim† Hae-Weon Lee† and Jong-Ho Lee*,†, § †

High-temperature Energy Materials Research Center, KIST, Seoul 02792, Korea

‡

Department of Materials Science & Engineering, Research Institute of Advanced Materials, Seoul National University, Seoul 08826, Korea §

Division of Nano & Information Technology, KIST School, University of Science and Technology, Seoul 02792, Korea

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ABSTRACT

Strain-induced fast ion conduction has been a research area of interest for nanoscale energy conversion and storage systems. However, because of significant discrepancies in the interpretation of strain effects, there remains a lack of understanding of how fast ionic transport can be achieved by strain effects and how strain can be controlled in a nanoscale system. In this study, we investigated strain effects on the ionic conductivity of Gd0.2Ce0.8O1.9-δ (100) thin films under well controlled experimental conditions, in which errors due to the external environment could not intervene during the conductivity measurement. In order to avoid any interference from perpendicular-to-surface defects, such as grain boundaries, the ionic conductivity was measured in the out-of-plane direction by electrochemical impedance spectroscopy analysis. With varying film thickness, we found that a thicker film has lower activation energy of ionic conduction. In addition, careful strain analysis using both reciprocal space mapping and strain mapping in transmission electron microscopy shows that a thicker film has higher tensile strain than a thinner film. Furthermore, the tensile strain of thicker film was mostly developed near a grain boundary, which indicates that intrinsic strain is dominant rather than epitaxial or thermal strain during thin film deposition and growth via the Volmer-Weber (island) growth mode.

KEYWORDS : gadolinium-doped ceria, strain effect, ionic conductivity, Volmer-Weber growth, intrinsic strain, reciprocal space mapping


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Fast oxygen-ion-conducting oxides have been widely used for catalysts,1 membranes,2 electrochemical sensors,3 and solid-oxide fuel cells (SOFCs)4 owing to their excellent kinetics properties involving oxygen ion migration. A conventional way to facilitate improving ionic conductivity of oxides is to substitute a host cation with an aliovalent-cation dopant, as in the case of gadolinium-doped ceria (GDC) and yttria-stabilized zirconia (YSZ). Doping with lower valent cation as an acceptor leads to an increase in oxygen vacancy concentration by following

the electro-neutrality condition: [ ] ≅ 2[ ∙∙ ].5–7 However, a method for improving ionic conductivity through dopant addition is usually limited not only by the solubility of the dopant cation, but also by the defect association between the oxygen vacancy and dopant cation, or between the oxygen vacancies at high doping level.8–10 Recently, further improvements in ionic conductivity than the aforementioned dopant engineering have often been achieved by epitaxial strain engineering, which enhances ionic conductivity via epitaxial strain induced by the lattice mismatch between a thin film and substrate or between multi-layered thin films.11–19 For example, Fluri et al.18 reported 0.35% epitaxial tensile strain lowers activation energy by ~ 0.05 eV in a samarium-doped ceria (SDC) thin film which is epitaxially grown on SrTiO3 (STO) buffer layer. Sanna et al.19 proposed a new design for chemically stable and fast ion conductors with the epitaxial strain in GDC/Er2O3stabilized δ-Bi2O3 (ESB) multi-layered thin films. In spite of these recent achievements, however, there still remains lack of understanding of how much fast ionic conduction can be achieved by the epitaxial strain due to many conflicting reports and lively debates regarding the influences of epitaxial strain on ionic conductivity.15,20–24 Especially, a colossal enhancement of ionic conductivity achieved by Barriocanal et al.12 in STO and YSZ multilayered thin film has hardly been verified in other studies, even those considering similar thin-film structures.20–24 3 ACS Paragon Plus Environment

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Until this time, the limitation of epitaxial strain engineering has been usually explained by three major factors. One is fast relaxation of epitaxial stain along thickness of thin film. Mostly, for oxide thin films, a critical thickness at which misfit dislocations are formed to relive the stress of thin-films, is expected at most to be around 10–20 nm 25–27 and thus it is hard to imagine that the epitaxial strain will remain in thicker film than the critical thickness. Lee et al.16 also described the relaxation profile of epitaxial strain over a relatively short distance range in GDC (111) epitaxial thin film. From the X-ray diffraction (XRD) observation and calculation of average strain as a function of thickness, ̃(t), which is expressed by

̃ =

(1)

where ε0 and α are constants, and x is the distance from the interface, they argued that 99% of strain can be released within 26 nm thickness. Another important but often ignored factor is the growth mode of thin film. As described in Scheme S1, thin films grown by the Volmer-Weber (island) growth mode which is the most common mechanism in heterostructured thin film growth18,28–31 go through dramatic change of strain. At the early stage of film growth, nucleation of discrete islands accompanies compressive strain due to high surface stress of islands. As the thin film grows and islands get closer, coalescence of islands occurs via a general “island zipping” process that minimizes the total energy of low mobility material by removing two surfaces and forming a grain boundary as a result of adhesive interaction between adjacent island edges,31 and thus induces tensile strain by elastic deformation around a grain boundary. After the deposited thin film has continuity, very complex situations involved in strain formation are intertwined to reach a steady state. For instance, grain growth and densification of thin film leads to tensile strain by eliminating excess4 ACS Paragon Plus Environment

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free volume of grain boundaries, whereas atomic migration from a free surface to a grain boundary to release tensile stress, or impingement of energetic particles is even attributed to tensile-to-compressive strain transition.28 In any case, under this island growth mode, the intrinsic strain is more dominant than the epitaxial strain through whole growth stages, because epitaxial strain can be released very easily by the formation of misfit dislocation at a heterointerface. Even in the case of multilayered-thin films, the intrinsic strain is also inevitable. For example, a layer-by-layer growth mode of A on B in an ABAB-multilayered thin film can be achieved when the surface energy of B is higher than sum of surface energy of A and interface energy (γB > γA + γAB).32 However, this growth mode can no longer be continued for the next layer growth (growth of B layer on top of A layer). In this case, the epitaxial thin film B can only grow by island growth mode. Despite this complex situation, however, most of the studies so far have only attempted to explain the conductivity enhancement phenomenon by the epitaxial strain effect rather than the intrinsic strain effect.18,30,33,34 The other major factor that has caused an interpretation error for the conductivity enhancement effect is the inadequate conductivity measurement that could not account for the influence of structural defects, such as grain or column boundaries. Since most thin films grown by the island growth mode have a columnar structure, the column boundary with a directionality perpendicular to the substrate will necessarily interfere with the ionic migration in a specific direction: in this case, the in-plane direction. Nonetheless, many studies have overlooked this perpendicular-tosurface defect and measured conductivity along the in-plane direction by only taking into account experimental convenience, thus, many studies have conflicting reports of strain effects and lacked reproducibility. For example, both Kant et al.34 and Fluri et al.18 observed the same enhancement phenomena of ionic conductivity from the in-plane conductivity measurement of 5 ACS Paragon Plus Environment

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epitaxial GDC and SDC thin film on an STO layer, where in-plane tensile strain is supposed to be induced by epitaxial correlation between ceria and STO. However, they reported a completely opposite tendency of the strain effect on activation energy of ionic conductivity. As to this opposite tendency, Kant et al. explained that structural defects to release the lattice strain might interfere with the activation energy and thus alter the strain effect on conductivity enhancement. In particular, when the perpendicular-to-surface defects, such as grain boundaries in a thin film are not controlled in an in-plane conductivity measurement, the measured in-plane conductivity must be regarded as distorted conduction, which is already affected by the perpendicular-tosurface defects.35 Hence, in order to avoid any interference from perpendicular-to-surface defects, ionic conductivity should be measured in the out-of-plane direction, although the experimental setup is difficult. In this study, through an out-of-plane conductivity measurement technique (Scheme S2), we investigate the strain-induced effect on out-of-plane ionic migration of (001)-oriented Gd0.2Ce0.8O1.9-δ

(GDC) epitaxial film deposited on a conductive substrate, single-crystalline

niobium-doped strontium titanate (Nb-STO). Epitaxial correlation between deposited film and substrate was confirmed by XRD analysis. In addition, cross-sectional transmission electron microscopy (TEM) images confirmed that GDC films are grown in highly-textured columnar structure with low-angle grain boundaries. (see Figure S2) Detailed discussions for thin-film epitaxy and microstructure are described in Method section in Supporting Information. Out-ofplane conductivity, which has no ionic conduction across the grain boundaries, was measured by electrochemical impedance spectroscopy analysis (EIS). Careful strain analysis was performed using both reciprocal space mapping (RSM) and strain mapping in TEM to analyze the local strain distribution within the films. Based on this local strain state analysis for films with various 6 ACS Paragon Plus Environment

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thicknesses, the exact mechanism for strain-induced fast ion conduction was identified and further insight to facilitate ionic conductivity in thin films by controlling the film’s strain state was obtained. Typical impedance spectra in Nyquist and Bode plots for out-of-plane conductivity measurements at 150 °C are shown in Figure 1(a) and (b). Since conductivity of Nb-STO substrate is much higher than GDC thin films ( σNb-STO ≅ 150 Ω,- cm,- at room temperature), the contribution of pure ohmic resistance is negligible and thus the impedance spectra generally contains two major components consisting of a parallel RQ element (GDC film) at high frequencies and a constant phase element (electrode reaction) at low frequencies. For out-of-plane conductivity measurements, grain boundary contributions are not commonly observed, owing to a parallel conduction path to the grain boundary (no conduction across the grain boundary), thus, the overall impedance spectra can be divided into two contributions: resistance to ionic migration through the grain of the thin film at high frequencies and the interface resistance at the interface between the thin film and electrode at a low frequencies.36–38 Therefore, the resistance and capacitance of GDC thin film can be estimated by fitting the impedance spectra using the equivalent circuit in the inset of Figure 1(a). The calculated capacitance of GDC thin film is in the rage of 3–70 nF, which corresponds to the value calculated using the relative permittivity of GDC36 and geometric factor for the out-of-plane direction (A/l ≅ 7 m). Moreover, the parallel RQ element in the high frequency range is independent of oxygen partial pressure and AC amplitude change, which clearly indicates that the high-frequency component is attributed to ionic migration of GDC thin films (Figure S5). The calculated ionic conductivity is plotted as a function of temperature in Figure 1(c). To compare the conductivity of GDC thin film with that of an unstrained bulk GDC, the ionic 7 ACS Paragon Plus Environment

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conductivity of the grain and the total (grain + grain boundary) conductivity of bulk GDC extrapolated from ref.[39] are added in Figure 1(c). As shown in Figure 1(c), the ionic conductivity of thin films is higher than the bulk sample by approximately two orders of magnitude, regardless of film thickness. Among them, the thinnest film (54 nm) shows the highest conductivity, and the conductivity tends to decrease with increasing film thickness. Interestingly, the out-of-plane conductivity of GDC thin films is close to that of grain in bulk GDC, which supports that grain boundary contribution can be effectively excluded via out-ofplane conductivity measurement. In addition, the activation energy for ionic migration calculated from

4567 =

° 89:;

=

exp A−

CD,F G=

H

(2)

° where T is temperature, R is the gas constant, 4567 is a pre-exponential factor, and IJ,K is the

free energy for oxygen vacancy migration, also varied with film thickness. As shown in Figure 1(d), the activation energy decreases with increasing film thickness. The interesting thing is that the activation energy of the thinnest film (54 nm), which shows greatest difference in ionic conductivity from the bulk, is most similar to that of bulk grain. The enhanced or suppressed ionic conductivity in a thin film is generally explained by the nano-size effects, such as grain boundaries,40,41 the space charge layer (SCL),42,43 and the residual strain at the hetero-interface44 because of their large volume contribution in a nanoscaled system. Among these three factors, however, ionic conduction through grain boundaries might contribute little in an epitaxial thin film, owing to low grain boundary density. Even the case where a thin film grows with columnar structure having columns or grain boundaries, the resistance of grain boundaries can be still excluded, at least in the out-of-plane direction, as 8 ACS Paragon Plus Environment

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mentioned in Figure 1(a). In addition, since the thickness of the SCL is related to the Debye length, which is inversely proportional to charge carrier concentration, the SCL effect can be also neglected for highly concentrated charge carrier systems, such as highly doped ceria (in this study: 20 mol% Gd-doped ceria).45,46 Therefore, the peculiar ionic conduction characteristics in the thin films shown in this study can be considered to be due to the residual strain effect in an epitaxial thin film. The reciprocal space maps (RSMs) of STO2L03 and GDC2L2L4 are scanned for three samples with different thicknesses, thinnest (54 nm), thickest (477 nm), and one of the medium thicknesses (139 nm), to quantify the residual strain in an epitaxial thin film. The RSMs, which are aligned by the asymmetric diffraction peak of STO2L03, are shown in Figure 2(a)–(c). In a reciprocal space, each in-plane and out-of-plane lattice constant can be found by a reciprocal of qx and qz, and can be compared to the reciprocal lattice point of unstrained-bulk GDC. As shown in Figure 2(a)–(c), from the point of unstrained-bulk GDC (marked as ★), all the peak position of thin films shifts to the direction in which the absolute value of qx decreases and qz increases with the increase of film thickness, which indicates that the lattice constant elongates for the inplane direction, while it decreases for the out-of-plane direction. Moreover, all RSMs show differences in peak broadening with film thickness. The peak broadening in RSMs is mainly explained by two factors. One is mosaic spread, which is caused by misorientation between grains. The mosaic spread makes the peaks scattered in ω-scan direction (perpendicular-toreciprocal lattice vector) due to variation of ω angle (∆ω). The other is grain boundaries, which are perpendicular to the surface. Grain boundaries contribute to peak broadening in the qx direction (parallel to surface); therefore, peak broadening in the qx direction (∆qx) is inversely proportional to the lateral coherent length (see Figure S5).47–50 9 ACS Paragon Plus Environment

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The out-of-plane lattice constant (aout), in-plane lattice constant (ain), lateral coherent length (∆qx−1), in-plane, and out-of-plane lattice strain (ε) calculated by

in, out =

Jin, out , Jbulk Jbulk

(3)

where abulk is lattice constant of unstrained bulk GDC, 5.423 Å (JCPDS card 75-0162), are listed in Table 1. As shown in Table 1, out-of-plane lattice constant decreases with increasing film thickness while in-plane lattice constant shows the opposite tendency. Therefore, to consider change in both out-of-plane and in-plane lattice constant, fractional lattice volume change (∆Vf) is calculated by ∆Vf = (Vfilm - Vbulk) / Vbulk where Vfilm = ain2aout, and Vbulk=abulk3, and included in Table 1. Noticeably, lattice volume expands with elongating in-plane lattice constant, which means migration volume for oxygen ion also can be dominated by in-plane tensile strain. In addition, a very unique tendency was observed in terms of lateral lattice constant and the strain state of the film with respect to film thickness. First, for lateral coherent length, the RSM peak of the thinnest 54-nm film clearly broadens in the ω-scan direction, which indicates mosaic spread is dominant in thinner film. However, peak broadening occurs in both the ω-scan and qx directions as the film thickness increases. As mentioned above, this might be attributed to high grain boundary density in thicker films, which can relax the high elastic energy; thus, a reciprocal of lateral broadening (∆qx−1) becomes two times larger in the thinnest 54-nm film. Nevertheless, the in-plane lattice constant of the thinnest 54-nm film was the smallest and its value was close to that of unstrained bulk, whereas the 477-nm thick film had the largest difference in lattice constant with bulk, thereby showing the largest residual tensile strain of 0.83%. This result is completely different from the conventional expectation that the strain state of the film has been generally interpreted by the epitaxial strain. If the overall strain state is 10 ACS Paragon Plus Environment

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governed by epitaxial strain, the lattice constant would become closer to that of bulk with increasing film thickness, owing to strain relaxation via defect formation, such as misfit dislocation. However, in our study, tensile strain was more likely to increase with increasing film thickness and decreasing lateral coherent length, which indicates that residual strain in thin films, especially those grown by the Volmer-Weber growth mode, can’t be interpreted solely by epitaxial strain. In order to estimate the thickness range affected by the epitaxial strain in our study, inter-planar lattice distance is analyzed from high-resolution-TEM image as a function of distance from the interface (Figure S3). According to the analysis for 54-nm and 139-nm thin film, in-plane tensile strain is mainly developed near film-substrate interface but rapidly relaxed within less than 20 nm thickness range. This results indicates that only small portion of film thickness is affected by the epitaxial stress (from rough estimation, ~10 % of thickness is affected by epitaxial stress in 139-nm thin film). Furthermore, the influence of epitaxial strain become further decreased as the film thickness increases. Continuing with Figure 1(d), the activation energy for ionic migration is plotted as a function of residual tensile strain in Figure 3, and also as a function of fractional lattice volume change in inset of Figure 3. As seen in the figure, a tensile strain of 0.83% given to the 477-nm thick film clearly lowers the activation energy by 0.05 eV, whereas the activation energy of nearly unstrained 54-nm film is close to that of unstrained bulk. The degree of reduction in activation energy in the ceria system by strain effects is also comparable with those of other studies. For example, according to Fluri et al.18, epitaxial (100) Sm-doped ceria thin film deposited on a BZO+STO buffer layer showed lower activation energy of in-plane ionic migration by 0.05 eV under 0.35% tensile strain. Since the overall lattice volume change appears to be dominated by increase of in-plane biaxial strain, the overall tendency and degree of reduction in activation 11 ACS Paragon Plus Environment

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energy might be similar to our study, in spite of their different measurement direction. In addition to the experimental results, Souza et al.51 reported the biaxial strain effect on ionic migration in un-doped ceria through a computational approach. According to their calculation results, the activation energy is expected to decrease by 0.1 eV for out-of-plane migration and 0.04 eV for in-plane migration under 0.8% tensile strain. Nevertheless, it is still questionable that the ionic conductivity of thicker films is smaller than that of thinner films in the whole measurement temperature range, in spite of its lower activation energy for ionic migration (Figure 1(c)). This can be partially attributed to grain boundary contributions to the reduction of effective conduction paths, because grain boundaries are still regarded as non-effective paths for ionic migration, even in columnar-structured film, owing to the SCL built around grain boundaries.39,45,46 Hence, if the non-effective path along a grain boundary and conductive path passing through the grain interior make a parallel circuit, the effective conduction area would obviously be reduced in the thicker film with higher grain boundary density to alleviate the accumulated elastic energy. For the in-depth analysis of microstructure and strain distribution of GDC thin films, 50-nm and 410-nm films were deposited on undoped STO (100) single-crystal substrates (MTI Corporation, USA) with the same deposition conditions as for the thin films deposited on NbSTO. HR-TEM images and inverse fast Fourier transform (IFFT) images of the (002) atomic plane are shown in Figure 4. For the 50-nm film, edge dislocation arrays are observed along lowangle boundaries, which indicates imperfection of coalescence between misoriented grains.35 From the distance of gaps between the edge dislocation arrays, the 50-nm film is expected to have a grain size of approximately 15–20 nm. On the other hand, for the 410-nm film, the misorientation angle seems larger than that of the 54-nm thin film, so that low-angle boundaries 12 ACS Paragon Plus Environment

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are clearly visible rather than edge dislocation in the IFFT image. Therefore, the grain size of the 410-nm film can be estimated to be approximately 5–10 nm from the IFFT image. Remarkably, both the estimated grain sizes from IFFT images are consistent with the lateral coherent length (∆qx−1) estimated from the RSMs. In addition to microstructural observations, the strain state of thin films was analyzed using a strain mapping system (TOPSPIN, NanoMEGAS, USA) in TEM and its spatial distribution within the thin film layer was analyzed by changing the selected area electron diffraction (SAED) pattern from a reference point; thus, a relative compressive or tensile strain was visualized on a bright-field TEM in pixel units. The bright-field images and strain maps of two samples, 50-nm and 410-nm films, are shown in Figure 5. The scale bar in Figure 5 indicates the relative strain in which compressive- and tensile-strain are expressed by minus and plus, respectively, and a reference point (marked by a white circle in Figure 5(a), (b)) is selected in the middle of the thin film for both samples. First, the bright-field images in Figure 5 clearly confirm the existence of grain boundaries for both 50-nm and 410-nm films as in other reports regarding island growth of ceria (100).18,52 Second, tensile strain is observed within a range of approximately 10 nm from both the interface and grain boundaries that are parallel and perpendicular to the substrate, respectively. The former strain might be attributed to epitaxial or thermal strain, owing to lattice and TEC mismatch between the film and substrate, whereas the latter might be attributed to intrinsic strain, as described in Scheme S1. However, the epitaxial and thermal strain near interface is likely to be fully released within 20 nm from the substrate, as expected from Figure S3 and previous studies.16,23,53 On the other hand, the extent to which the intrinsic strain (residual tensile strain near grain boundaries) may affect the overall strain state, depends on the concentration of grain boundaries. For example, if the concentration of grain boundaries is 13 ACS Paragon Plus Environment

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sufficiently large, the strain state of the entire film will be subjected to a tensile strain state. Indeed, as shown in Figure 5, tensile strain is uniformly distributed throughout the whole 410-nm film owing to high grain boundary density, whereas it is less in the 50-nm thin film. Consequently, among the three aforementioned strain sources, intrinsic strain can be regarded as the main factor for GDC epitaxial thin film, because the epitaxial strain is fully relaxed beyond a few-nm range from the interface, owing to high grain boundary density. In addition, only approximately 0.1% of tensile strain is expected by a TEC mismatch, which is calculated by TEC of bulk GDC and STO.54,55 Solovyov et al.50 also reported the correlation between lateral grain size and intrinsic tensile strain, which is induced by joining of Volmer-Weber islands. According to them, the average lateral grain size is inversely proportional to lateral strain in ceria (100) epitaxial thin film, which is in accordance with our observed correlation between intrinsic tensile strain and lateral coherent length. One more observation in their study is that the intrinsic tensile strain is fairly stable and can only be released by grain growth through high-temperature annealing at above 1100 °C, which is much higher than general thin film deposition temperatures. Although general stress evolution during island growth mode follows typical compressivetensile-relaxed (or compressive) stress behavior, a large tensile stress has been frequently observed in low-mobility materials, such as ionic solids,50 low-mobility metals,56 or metal thin films deposited at low temperature28 owing to sluggish relaxation kinetics. The level of intrinsic tensile stress contribution to residual strain of a thin film can be estimated by Hoffman’s model.57 Hoffman described coalescence as a process to compensate surface energy with a onedimensional elastic deformation when the gap of adjacent two islands reaches a critical size and defined the intrinsic stress by coalescence (σint) as


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457 = U

VWXY ,XZ[ \

(4)

where Y is Young’s modulus, γs and γgb are the surface and grain boundary energy, respectively, and 2r is the grain diameter. Based on Eq. (4), the intrinsic stress of our 410-nm GDC film was roughly estimated using literature data of Y = 187 GPa and 2]^ − ]_` ≅ 0.2 ~ 2.0 J/cm2,58,59 and the experimental result of 2r = 10 nm from the TEM image in Figure 5. Interestingly, the estimated intrinsic stress is fairly large, in the range of 2.7–8.6 GPa, compared to other reported levels.18,60 Furthermore, the intrinsic strain (εint) level induced by this range of intrinsic stress is expected to be 1.1 to 3.2% from the biaxial stress-strain relationship, σ = Yε/(1 − ν), where ν is the Poisson’s ratio (νGDC = 0.3061), which is comparable to both the theoretical epitaxial strain of 1.8% in GDC/STO heterostructures and the experimentally observed strain of 0.8% in this work. In conclusion, the intrinsic strain, which has been frequently neglected in the studies of straininduced fast-ion-conduction in oxide electrolyte should always be considered, especially for thin films grown by the Volmer-Weber growth mode. Further investigation to identify how large intrinsic-tensile strain can be produced by coalescence and how it can reach maximum strain by controlling coalescence in Volmer-Weber growth would be treated in a forthcoming study. We investigated out-of-plane ionic transport in strained GDC thin films by using conductive single-crystal Nb-STO substrates. The out-of-plane conductivity measurements allow the strain effects on ionic conductivity to be observed without interfering with perpendicular-to-surface defects, such as grain/column boundaries that have been overlooked until now. From the extensive microstructure analysis with RSM and TEM, we found that large tensile strain is induced in an epitaxial thin film grown by the Volmer-Weber growth mode. Remarkably, for local strain distribution of thin films, it is noticeable that larger tensile strain is developed for the 15 ACS Paragon Plus Environment

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film with finer grain size and formed intensively near grain boundaries rather than the film/substrate interface, which implies that the residual strain of GDC thin films is attributed to intrinsic stress by coalescence. Furthermore, it is also highlighted that the tensile strain of 0.83% induced by intrinsic stress indeed lowers the activation energy for oxygen-ion conduction through grain interior by ~0.05 eV, which is consistent with the results described in other computational and experimental studies. To the best of our knowledge, this is the first experimental evidence to show the significance of the intrinsic strain effect on ionic conductivity. Unlike the epitaxial and thermal strain that is confined to a small spatial range under specific conditions, this intrinsic strain can be developed extensively, regardless of the type of film or substrate materials and the thickness of the film. Therefore, we believe that this study can shed light on effectively utilizing the strain effect for the performance improvement of energy conversion and storage devices based on ion conducting thin films, such as µ-SOFCs, oxygen permeation membranes, and sensors.


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FIGURES

Figure 1. (a) Nyquist and (b) bode plots for out-of-plane conductivity measurements with varying film thickness, chi-square of 0.01~0.001 indicate the equivalent circuit well describes the impedance spectra, (c) Arrhenius plot for out-of-plane ionic conductivity of GDC thin films; solid line indicates total ionic conductivity (σgrain + σg.b.) and dashed line indicates σgrain from ref.[39], (d) activation energy for ionic migration with varying film thickness; dashed line indicates the activation energy of σgrain obtained from ref.[39]


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Figure 2. RSMs of (a) 477 nm (b) 139 nm (c) 54 nm GDC thin films; the arrow in (c) indicates a dc de), the reciprocal lattice point of unstrained-bulk GDC is marked reciprocal lattice vector of (c as a star. 18 ACS Paragon Plus Environment

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Figure 3. Activation energy (Ea,m) for ionic migration as a function of in-plane strain (εin); dashed line indicates activation energy of σgrain from ref.[39], activation energy is also plotted as a function of fractional lattice volume change (∆Vf) in inset to considering change of both inplane and out-of-plane lattice constant.

Figure 4. HR-TEM and IFFT of (002) images for (a), (b) 50-nm and (d), (e) 410-nm thin film, (c), (f) are magnified images of (b), (e); for 50-nm thin film, edge dislocations are marked as red symbols, and visible low-angle boundaries are marked as red dashed lines for 410-nm thin film. 19 ACS Paragon Plus Environment

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Figure 5. Strain mapping (SM) and cross-sectional virtual bright-field (VBF) images of the mapping area for (a) 50-nm and (b) 410-nm thin film; a white circle in SM images is used as a reference point of SAED pattern; some specific area in which the grain boundaries exist is indicated by red-dot circles in VBF images.

ain / Å

aout / Å

∆Vf / %

∆qx-1 / Å

εin / %

εout / %

477 nm

5.468

5.407

1.37

92.38

0.83

-0.29

139 nm

5.457

5.420

1.20

147.84

0.63

-0.05

54 nm

5.423

5.431

0.15

211.37

-

0.15

Table 1. Out-of-plane lattice constant (aout), in-plane lattice constant (ain), fractional lattice volume change (∆Vf), lateral coherent length (∆qx−1), in-plane strain (εin), and out-of-plane strain (εout) of GDC thin films.


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ASSOCIATED CONTENT Supporting Information. Detailed information for method including the crystal structure of thin films, microstructure of thin films, thin-film conductivity measurements, the epitaxial strain analysis, thin-film composition, and reciprocal space mapping (PDF)

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]. Author Contributions J.A. †, ‡, H.W.J. ‡, and J.H.L.†, § planned this study and co-wrote the manuscript. J.A. †, ‡ conducted the experiments and characterizations. K.J.Y. †, H.K. †, and B.K.K. † advised on the results regarding electrochemical measurements. H.J. †, J.W.S. †, §, and H.W.L. † advised on the results regarding the physical properties. All authors have given approval to the final version of the manuscript. ACKNOWLEDGMENT This work was supported by the Technology Development Program to Solve Climate Changes of the National Research Foundation (NRF) grant funded by the Korea government (Ministry of Science and ICT) (NRF-2017M1A2A2044982).


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