The Influence of Local Distortions on Proton Mobility in Acceptor

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The Influence of Local Distortions on Proton Mobility in Acceptor Doped Perovskites Jilai Ding,†,‡,∇ Janakiraman Balachandran,‡,∇ Xiahan Sang,‡ Wei Guo,‡ Jonathan S. Anchell,‡ Gabriel M. Veith,§ Craig A. Bridges,∥ Yongqiang Cheng,⊥ Christopher M. Rouleau,‡ Jonathan D. Poplawsky,‡ Nazanin Bassiri-Gharb,†,# Raymond R. Unocic,*,‡ and P. Ganesh*,‡ †

Department of Materials Science and Engineering, Georgia Institute of Technology, Atlanta, Georgia 30332, United States Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States § Material Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States ∥ Chemical Sciences Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States ⊥ Spallation Neutron Science, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States # The George Woodruff School of Mechanical Engineering, Georgia Institute of Technology, Atlanta, Georgia 30332, United States

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‡

S Supporting Information *

ABSTRACT: Optimizing proton conduction in solids remains the most promising solution for achieving intermediate temperature (∼750−1000 K) solid oxide fuel cell devices, and enabling selective membranes for H2 separation. Proton conduction, a thermally activated process, exhibits its highest rates in yttrium (Y) acceptor doped BaZrO3 at an optimal doping level of 20% Y. The presence of extended defects such as grain boundaries has typically generated a wide variability in reported conductivity values. This has hindered a fundamental mechanistic understanding of how (acceptor) doping levels correlate with the activation energy of protons to produce an optimal doping level for fast proton transport. While isolated dopants have been suggested as the primary source of proton trapping, our results indicate that it is the local dopantdensity that matters. Here, we show that increasing the local dopant density promotes localized lattice distortions in the presence of point defects such as oxygen-vacancies or proton interstitials. An increasing distortion amplitude traps the point defects more strongly in the form of polarons, forming defect-clusters at higher concentrations. This leads to a monotonic increase in the activation energy (and hence a decrease in proton mobility) as observed in our measurements. The optimum doping level can now be explained as a competition between increasing proton concentration with doping levels and increasing activation energy due to defect-clusters formed by defect-polarons. Based on our findings, we demonstrate how to improve proton conductivity in doped BaZrO3, by inhibiting this dopant-lattice polaronic interaction. This approach should be generally applicable for ionic conduction in perovskite oxides such as oxygen-ion conduction in solid-oxide fuel cells and alkali-ion conduction in solid-state batteries where carriers might get trapped as defect-polarons.

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INTRODUCTION

(BZO), prepared by pulsed laser deposition, have the cubic ABO3 perovskite structure, and have shown the highest conductivity of all known materials, with an Ea ∼ 0.63 eV for proton transport.2 Here stoichiometric refers to a 1:1 ratio of Ba:(Zr/Y) cations. Extended defects such as grain boundaries and dislocations further increase the activation energy, and as such polycrystalline samples show an activation energy higher than 0.7 eV.2−7 Beyond these microstructural effects, it is widely accepted that protons bind more strongly to the acceptor dopant, due to its intrinsic negative charge, and this binding can adversely affect proton conductivity.8,9

Proton conducting oxides are a promising class of materials for intermediate temperature solid oxide fuel-cells, due to the lower expected activation energy for proton conduction compared to oxygen ions in traditional solid oxide fuel cells.1 The ionic transport is quantified by ionic conductivity (σ) which is written as σ = nzeμ, where “n” is the density of free carriers with charge “ze” and mobility μ = μ0exp(−Ea/kBT), where Ea is the activation energy for the carriers to hop. Protonic carriers are introduced in the conducting oxide by filling the intrinsic oxygen vacancies that get created to compensate for extrinsic acceptor substitutional dopants. They diffuse through the material when thermally activated. Increasing activation energy implies a decreasing mobility. Stoichiometric thin films of 20% yttrium (Y3+) doped BaZrO3 © XXXX American Chemical Society

Received: February 2, 2018 Revised: June 22, 2018 Published: June 27, 2018 A

DOI: 10.1021/acs.chemmater.8b00502 Chem. Mater. XXXX, XXX, XXX−XXX

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molecular vibrations in either crystalline or disordered materials over a broad energy range (5- 600 meV). The INS experiment was conducted at a temperature of 5 K. STEM Measurements and Distortion Analysis. An aberration corrected Nion UltraSTEM 100 (operated at 100 kV) was used to obtain atomic resolution images of the Y-BZO specimens. High-angle annular dark field (HAADF) scanning transmission electron microscopy (STEM) images were acquired with an electron probe convergence semiangle of 31 mrad and 86−200 mrad collection semiangle. The STEM images were averaged from 10 fast frames acquired with a pixel dwell time of 4 μs and 512 × 512 frame size. The frame averaging process significantly reduces scan distortion and enhances signal-to-noise ratio. The locations of atomic columns were identified using normalized cross-correlation and fitted using a 2D Gaussian distribution. Two types of distortions, namely bond angle and displacement were calculated for each B-site atom, and were performed on at least three images for each dopant concentration. Bond angle deviation is calculated by the difference between each B-site bond angle and 90 deg, showing how the lattice deviates from the cubic structure. Displacement is defined to be the distance between the observed B-site atoms and the unit cell center averaged from its four nearest A-site atoms. For 20Y-BZO, the distortions calculated from STEM images are not as high as those from DFT, which is due to the averaging effect over sample thickness, and a wide range of local environments spatially, not captured in the single supercell models used in the DFT calculations. The original STEM images without distortion overlays and distortion analysis on 5, 10 and 15Y-BZO are shown in SI Figure S6. tr-KPFM Measurements and Data Analysis. Energy discovery platforms were used in this study, which combines in situ surface potential mapping by tr-KPFM and nanofabricated thin film lateral devices. The tr-KPFM measurements were performed on a Bruker multimode Atomic Force Microscopy (AFM) equipped with a Nanonis controller. Cantilevers with conductive Cr/Pt coated tips (Budget Sensors, Co.; reasonance frequency ∼75 kHz) were used. During measurements, 15 V of DC bias and 1 V of AC bias (ω = 44 kHz) were applied to the cantilever. An external function generator (DS345, Stanford Research) was used to apply the AC waveform to the tip, and an external lock-in amplifier (SR844, Stanford Research) was used for signal processing. For tr-KPFM measurements, the tip scanned over spatially defined grid points of at least 30 (between electrodes) × 3 (along electrodes). At each point, a time-resolved two-step measurement was performed: for the first t seconds (t ranges from 2 to 100 s, depending on the time constant, or the potential changing rate at a certain temperature) 30 V DC bias was applied through one electrode (biased electrode), and the other electrode was grounded (grounded electrode); and for the following t seconds, both electrodes were grounded to allow the accumulated charge carriers to relax. Some of the outlier points in the mapsdue to the abrupt changes in topography (scanning tip encountering surface particle or electrode edge)were removed manually. Measurements were performed in a temperature range from 25 to 200 °C, and relative humidity of 90% (as measured at room temperature) using a gas cell. The noncontact AFM tip scans over the spatially defined grid points spanning the two Pt electrodes, measuring surface potential variation as a function of time while an externally applied DC bias is applied through the lateral electrodes. The obtained surface potential mappings are plotted as a function of different temperatures and dopant concentrations, shown in SI Figure S3. The surface potential changes can be fitted to an exponential model, Φ = A + B · e−t/τ, where the time constant τ is inversely proportional to the transport process of protons. The time constant is obtained for each sample as a function of temperature (SI Figure S2), and the activation energy is obtained through an Arrhenius plot of each sample, as shown in Figure 1a. Atom-Probe-Tomography (APT) Measurements and Data Analysis. An FEI Nova 200 dual-beam focused ion beam (FIB) instrument was used to perform lift-outs and annular milling to fabricate the needle-shaped APT specimens containing the Y-BZO layer. A wedge lift-out geometry was used to mount multiple samples

Calorimetry measurements on powder samples have indicated that defect-clusters form at high dopant concentrations,10 but we do not know their nature. It has also been speculated that localized bound polarons11−15 might be present in acceptordoped oxides. But, as pointed out by K. D. Kreuer,1,16 the real correlation among spatial distribution of dopants as a function of doping level, the nature of defect-clusters that might form, and how they affect proton binding and mobility through the material is subject to debate. Recently, it has been shown that stoichiometric 20% Ydoped BZO thin films exhibit high proton conductivity at operating temperatures near 800 K.2 We similarly grow epitaxial and polycrystalline Y-BZO thin films by pulsed laser deposition (PLD) on (100) oriented MgO substrates (Supporting Information (SI) Figure S1). In addition to novel synthesis, we have performed a range of complementary experiments (neutron spectroscopy, electron microscopy, atom-probe tomography and transport measurements), guided by first-principles theory in every step of the way, to systematically investigate Y-doped BaZrO3 across the 0−20% range of doping levels (0, 5, 10, 15, and 20Y-BZO), and comprehensively elucidate the localized interactions that protons experience around dopant chemistry, how it depends on the dopant concentration and influences proton mobility and find ways to rationally improve proton transport in this class of materials. This deeper atomistic understanding not only rationalizes seemingly disparate key recent experimental findings in the literature2,15,17 but also allows us to propose a rational design principle to improve ionic conductivity in perovskite oxides such as oxygen-ion conduction in solid-oxide fuel cells18 and alkali−ion conduction in solid-state batteries19 and has serious implications in far-reaching novel applications of quantum perovskites such as fuel-cell,20 pH sensing21 and artificial-synapsis.22

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EXPERIMENTAL AND THEORETICAL METHODS

Sample Preparation. In this study, great effort has been made to prepare a dense PLD target with a variety of dopant concentrations, and to determine deposition parameters for smooth, dense and epitaxial thin films. 0%, 5%, 10%, 15%, and 20% mole percent Y-BZO pellets were prepared by powder sintering at 1500 °C for 24 h. Both epitaxial and polycrystalline Y-BZO films (∼500 nm) were created on oriented MgO substrates by pulsed laser deposition (PLD). A KrF excimer laser (Coherent Lambda Physik GmbH) with a wavelength of 248 nm and a pulse width of 25 ns was focused on the target with a spot area of about 2 mm2. The epitaxial films were deposited at 750 °C while the polycrystalline films were deposited at 600 °C. The target-to-substrate distance was 40 mm, and the oxygen partial pressure was 40 mTorr. The thermal contact between the sample holder and the deposition substrate was provided by Ag paste. The laser energy density was about 1.5 J/cm2 with a repetition rate of 10 Hz. The growth rate of the films was about 3 Å/s. The PLD system was equipped with a reflection high-energy electron diffraction (RHEED) system for the in situ diagnostic of the deposition process. The out-of-plane crystal structure of Y-BZO samples were checked by a Philips Xpert X-ray diffractometer (XRD, Cu Kα = 1.5418 Å). For time-resolved Kelvin probe force microscopy (tr-KPFM) measurements, lateral Cr/Pt electrodes (20 nm/80 nm thick) were subsequently created by evaporating and patterned by a lift-off method. The electrodes were 10 μm wide with interelectrode distance of 70 μm. Neutron Scattering and Modeling. The vibrational (phonon) spectrum of the wet and the dry samples were experimentally determined through inelastic neutron scattering (INS) employing the VISION spectrometer at the Spallation Neutron Source in ORNL. VISION is an inverted geometry instrument that can characterize B

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We would like to note that while the trends in atomic distortions and formation- and migration-energies due to point-defects and defectcomplexes are readily captured using our semilocal DFT functional, the localization of the electron or hole associated with any boundstate (such as a bound hole-polaron) requires self-interaction corrections and can be dependent on the choice of the functional and its parameters to make this correction.14 Indeed, the PBE-DFT calculated vibrational density-of-states for cubic BaZrO3 is in excellent agreement with INS measurements, as shown in SI Figure S5. Not only are the bulk modes in agreement with experiments, but simulations also correctly predict the dependence of the oxygenproton bond strengths on the local B-site atomic environment in Ydoped BZO. In our recent first-principles based screening approach we were able to obtain trends in the formation energy of point-defect (and complexes) using the PBE-functional, in agreement with the more accurate hybrid-functional calculations.25

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RESULTS AND DISCUSSION Proton activation energies from our tr-KPFM measurements26 (see SI for details on how transport analysis was performed using tr-KPFM method) on the epitaxial thin films is shown in Figure 1a (Epi Y-BZO). We compare the obtained activation energy for proton transport with all earlier measurements, typically in polycrystalline Y-BZO samples.2,6,7,27−30 As seen, these results from the literature have wide variability, and hence cannot be used to understand the dependence of the intrinsic proton mobility on dopant concentrations. In comparison, the only thin-film stoichiometric Y-BZO electrochemical impedance spectroscopy measurement we found in the literature at 20% Y-BZO,2 has an activation energy of 0.63 eV, comparing very well with our measured activation energy of 0.62 eV at this same composition, suggesting that thin-films are good model systems for our purposes. These results also suggest that tr-KPFM measurements can be used as a benchmark to determine the evolution of the activation energy at lower levels of doping (also see SI Figures S2 and S3). A monotonic decrease of activation energy is observed with decreasing dopant levels. The lowest activation energy appears to be at 5% Y-doping, and is ∼0.45 eV, appreciably lower than the value measured at 20% yttrium concentration. This is a surprising result, and points to an intrinsic effect on the activation energy solely due to dopant distribution in the host BZO material. The fact that protons prefer to adsorb near acceptor dopants cannot alone explain this monotonic trend.9 Polycrystalline samples show a similar surprising trend (Poly Y- BZO in Figure 1(a)) and as expected, it has an activation energy higher than epitaxial samples. It is unclear what causes the activation energy to increase with dopant concentration in epitaxial Y-BZO thin films, even without any extended defects. Yttrium-cluster based defectcluster formation, due to solubility limits, has been speculated10 to explain the decrease in mobility observed in bulk samples with increasing yttrium concentration. However, such yttrium-clusters have neither been experimentally observed, nor is it clear how their existence would result in modified mobility of protons. Figure 1b depicts the spatial distribution of Y atoms in a 3D volume for 20Y-BZO sample from atom probe tomography (APT) measurements31 (see SI for details on how yttrium distribution was analyzed using APT data in 5- and 20Y-BZO samples), suggesting that even at 20% doping, Y atoms are fully miscible in BZO and are homogeneously distributed across the sample. However, even for this homogeneous distribution of Y dopants, an increase in the dopant concentration will increase the local dopant

Figure 1. (a) A summary of the activation energy for proton conduction in Y-BZO as a function of Y concentration in this study (red for epitaxial and orange for polycrystalline thin film samples) and other studies; (b) A reconstructed APT volume showing the homogeneous distribution of yttrium atoms (yellowish green) for 20Y-BZO sample; (c) Normalized distribution of the first Y distance around every Y atom as a function of dopant concentration. on a Si microchip coupon to enable the fabrication of multiple needles from one wedge lift-out. APT was performed with a CAMECA instruments LEAP 4000X HR (∼36% detection efficiency) and a 5000 XS (∼80% detection efficiency) local electrode atom probe. Samples were run at a base temperature of 40 K, applying 355 nm wavelength 10 ps laser pulses of 50 pJ at a repetition rate of 200 kHz. The data sets were reconstructed and analyzed using the IVAS 3.6.12 software (CAMECA Instruments). Through successive evaporation of the material, APT allows identification of the location of each atom not only of the surface, but also the bulk. First-Principles Theory. We perform first-principles calculations based on density functional theory (DFT) to compute the defect formation energies and defect migration energies employing Vienna Ab initio Simulation Package (VASP) with the generalized gradient approximation (GGA) according to Perdew, Burke, and Ernzerhof and the projector augmented-wave method.23 All the calculations on a 3 × 3 × 3 super cell encompassing 135 atoms. A Γ point centered 2 × 2 × 2 k-point mesh was used for Brillouin zone integration and a plane wave energy cut off of 400 eV was used in all calculations. The lattice constant was chosen to represent the experimental value (4.197 Å). The structural relaxations were performed to relax the internal coordinates without any change to the volume until the residual force of each atom was less than 0.01 eV/Å. The fully relaxed static transition energy barrier was obtained employing the nudged elastic band method (NEB) where the residue NEB forces on the atoms were less than 0.05 eV/Å. The ab initio molecular dynamics simulations24 were performed for a canonical ensemble employing the Nose algorithm. We employ a time step of about 0.5 fs and the simulation is run for about ∼12 ps. Data analysis is performed after an initial transient of 5−6 ps. Maximum displacement of the oxygen atom coordinating the dopant is plotted as max|ΔRO| in Figure 4, and will act as a proxy to the strength of the oxygen sublattice distortion. C

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obtain at 0K for a proton trapped between first neighbor Y atoms (i.e., Y2(1NN) model). In addition to agreeing with Yamazaki et al.’s study at 20% dopant concentration, our calculations (Figure 2(a)) clearly show that the proton-dopant interaction energy is not a constant for Y-BZO, but should have a concentration dependence. This should in turn lead to a dopant concentration dependent proton activation energy as measured by our tr-KPFM results (Figure 1a). Our theoretical calculations, validated against neutron diffraction measurements (SI Figure S6), also suggest that this concentration dependence could be related to an increasing local distortion with increasing dopant concentration. To experimentally confirm this theoretical prediction, we performed atomic resolution scanning transmission electron microscopy (STEM) to image the Y-BZO epitaxial thin films at different dopant concentrations (SI Figure S8 and S9). We measured local distortions by performing two types of distortion analysis of the local B-site coordination polyhedral directly from STEM. The spatial distribution of both types of distortions for BZO and 20Y-BZO34 (as shown in Figure 3a, b, e, f) supports our previous conclusions from APT measurements of random dopant distribution in our thin films. Histograms of these distortions (Figure 3c, g) show a shift to higher values of both its peak position and standard deviation with increase in dopant concentration. This implies that there is both an increase in the strength as well as the variability of the lattice distortion. The corresponding simulated histograms obtained from DFT calculations (Figure 3d, h) show a similar trend (also see SI Figure S10 for corresponding histograms for dry samples). The monotonic increase in the lattice distortion with increasing dopant concentration is due to a need to electrostrictively32 screen the increasing charges from the point defects (and their complexes) with increasing dopant concentration. The increasing distortion correlates with an increase in our measured activation energy for protons (Figure 1a). While a recent neutron study claims existence of protonpolarons in a related proton conducting cubic perovskite materials (Y-doped BaCeO3)15 we present a visual proof of the existence of such localized polaronic distortions in proton conducting oxides, and also discover that the strength of these distortions are dependent on the dopant concentration. We computationally tested the generality of this correlation between the lattice distortion and the activation energy to other dopant elements. As a proxy to the activation energy, we compute ΔEiMZr′,Hi· and correlate it with the maximum displacement of the oxygen atom coordinating the dopant (max|ΔRO|), for all possible dopants. As shown in Figure 4a, there is a “weak” volcano-like behavior of ΔEiMZr′,Hi· with max|ΔRO| at smaller values of the distortion, but with increasing max|ΔRO| the dopant-proton interaction energy generally becomes more favorable, indicating an overall positive correlation between oxygen distortion and proton trapping. The change in ΔEiMZr′,Hi· due to the local distortion, that is, the binding strength of the proton-polaron − δHi−ph, is approximately −1.20 eV for Y-BZO. This is a large value that drives the dopant-proton interactions from being repulsive ( > 0) to attractive (ΔEiMZr′,Hi· < 0). δHi−ph also correlates positively with (max|ΔRO|) as shown in Figure 4b. This suggests that there is perhaps some property of the dopant element that can act as a descriptor for the activation energy.

density. This results in an increase in the number of first nearest neighbor (1NN) Y−Y pairs, a measure of the strength of Y−Y association, as simulated in Figure 1c. In order to link increased dopant−dopant association with our observed increase in the activation energy for proton conduction, theoretical studies were performed. We performed DFT calculations of dopant (YZr ′ ) - interstitial (OHO· ) i · interaction energies (ΔE MZr′,Hi ) for four different dopant models with varying nearest neighbor (NN) separations of dopant atoms (see SI Figure S11 for details). This allows us to understand how dopant−dopant association affects formation of defects, such as oxygen vacancies or proton interstitials, thereby influencing proton conduction. As it is known,8,9 our neutron scattering experiments (see SI Figure S5 for details) suggest that the proton prefers to bind next to an yttrium atom (Y1 model). But surprisingly, as we increase the local density of dopant atoms around the defect, the interaction energy goes up, and is highest for the Y2 (1NN) model (Figure 2a). We

Figure 2. (a) The dopant-proton interaction energy (representing protonated samples), the corresponding maximum oxygen displacement and maximum proton transition energy among the various transition pathways for different model systems; (b) A schematic showing the correlation between proton binding energy and anionic (O2−) sublattice distortion for various local environments such as pure BZO, in one yttrium dopant environment (Y1) and in two yttrium (1NN) dopant environment Y2(1NN).

also find that the local atomic distortion, measured as the maximum displacement of sublattice oxygen atom (max|ΔRO|), directly correlates with it. Similar trend was observed for oxygen vacancy defects as well (see SI Figure S7). The oxygen sublattice distortion is possibly an electrostrictive effect and points to the formation of a localized phonon mode that condenses in response to a charged impurity, akin to formation of a proton-polaron.11,15,32,33 The Y2 (1NN) model also shows the highest migration barrier for protons (Figure 2a). This suggests that an increase in the number of Y−Y neighbors, can lead to an increase in the trapping of defects, due to the increased local distortion, thereby leading to an increase in the activation energy, as schematically shown in Figure 2b. Indeed, the study by Yamazaki et al.9 in 20% doped Y-BZO measured a proton-dopant association (or interaction) energy of −0.301 eV (−29 kJ/mol), close to the ΔEiMZr′,Hi· = −0.25 eV value we D

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Figure 4. (a) The dopant-proton interaction energy (ΔEiMZr′,Hi·) plotted as a function of maximum oxygen displacement. (b) The proton-polaron binding strength (δHi−ph) plotted as a function of maximum oxygen displacement. (c) Maximum oxygen displacement (max|ΔRO|), in a protonated model system as a function of the ratio of dopant ionic radius to the host cation ionic radius. (d) Mean squared displacement (MSD) of proton interstitials obtained from ab initio molecular dynamics (AIMD) simulations for various model systems. (A-site dopants are labeled in red, and B-site dopants in black).

If all B-site acceptor dopants in BZO create sublattice polaronic distortions and trap protons, then what is a good design route to improve conductivity? Our results suggest that by decreasing the polaronic lattice distortion, it might be possible to reduce proton trapping, and thereby lower the activation energy, increasing its mobility. To support this hypothesis, we performed DFT calculations on all possible acceptor dopants in the ionically bonded A-site Ba ions. For potassium (K) as the A-site dopant, the max|ΔRO| in this system is about 0.15 Å, which is much lower compared to the B site acceptor dopants and results in weaker ΔEiMZr′,Hi· and δHi−ph as shown in Figure 4a and 4b. Interestingly, the radii ratio of K is close to 1.0 and nicely falls on the observed trend of the distortion with the radii ratio in Figure 4c. This hypothesis is supported by our AIMD simulation on K doped BZO at ∼900 K showing a higher mean squared displacement (MSD) value compared to the protons trapped near Y dopant atoms as shown in (Figure 4d). Protons still appear to be fastest away from any dopant atom in BZO as seen from their MSD values (SI Figure S11 and S12 shows the proton trajectories). While our results qualitatively agree with the recent DFT calculations by A. Loken et al.37 that doping A-site atoms in BZO can lead to improved proton mobility, our current analysis of trends across all possible A/B-site dopant atoms clearly explains the reason for this -- doping A-site, especially with dopant atoms with a radius-ratio ≥1, leads to significant lowering of local distortions even at very high dopant concentrations, thereby improving proton mobility. This clear understanding can be used to rationally improve ionic mobility for any ionic species in a solid.

Figure 3. (a,b) Atomic resolution STEM imaging and data analytics to identify the bond angle deviation mapping for (a)0Y-BZO samples and (b) 20Y-BZO samples. (c) Bond angle deviation distribution as a function of dopant concentration obtained from STEM data. (d) DFT simulated bond angle deviation distribution for dry model samples; (e,f) STEM image data identifying displacement mapping for (e) 0Y-BZO samples and (f) 20Y-BZO samples. (g) Displacement distribution as a function of dopant concentration obtained from STEM data. (h) DFT simulated displacement distribution for hydrated model samples.

We found that the ratio of the radii of dopant atom with that of A/B-site atom can be such a descriptor. Indeed, as shown in Figure 4c the max|ΔRO| shows a generally increasing trend with increasing radii ratio. This finding reconciles the volcanolike plot of conductivity with dopant radii as measured in doped−BZO35 and observed in earlier simulation studies of doped−BZO8 and SrCeO336 with our observations of a “weak” volcano-like behavior of the proton-dopant interaction energy with the amplitude of local distortion as seen in Figure 4a, and points to the negative influence of local polaronic distortions on fast proton migration as recently seen in experiments.15 Generally, having small local distortions appears to be favorable for faster proton migration, and thus represents a useful design variable for proton conductors, alongside other factors influencing the overall proton conductivity such as dopant solubility, sinterability, and grain-size.

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CONCLUSIONS To summarize, we discovered that as the local dopant density increases with increasing yttrium concentration, it becomes E

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actively involved in interpreting the results and writing the manuscript.

easier to condense a localized polaronic lattice distortion to trap protons, especially around the monotonically increasing number of Y−Y pairs, thereby leading to a monotonic decrease in its mobility. Our findings go one crucial step beyond our current understanding from literature that protons get trapped near Y-dopants in Y-doped BZO, thereby explaining trends in mobility with dopant concentration. This is an important breakthrough. As we showed, this discovery is generally true for other dopant elements, and that the dopant-proton interaction energy becomes more favorable with increasing local lattice distortions. Based on this atomistic understanding of what limits proton mobilities in doped BaZrO3, we proposed a new route to design improved proton conducting material systems by inhibiting this dopant-lattice interaction. We expect that this adverse effect of dopant-lattice polaronic interaction on ionic conductivity might be present in many other energy storage systems, such as oxygen ion conducting solid-oxide fuel cells18 or alkali-ion conducting solid-state battery materials,19 which can be removed to potentially improve their performance dramatically. Our current findings have far reaching consequences in understanding some of the charge-lattice interaction mediated effects in quantum perovskite family of materials for fuel-cell,20 pH sensing21 and artificial-synaptic22 applications.

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Funding

This research was sponsored by the Laboratory Directed Research and Development Program (LDRD) of Oak Ridge National Laboratory, managed by UT-Battelle, LLC, for the U.S. Department of Energy (Project ID 7448). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Part of this research used resources of the National Energy Research Scientific Computing Center, which is supported by the Office of Science of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. Part of this research used resources of the Oak Ridge Leadership Computing Facility at the Oak Ridge National Laboratory, which is supported by the Office of Science of the U.S. Department of Energy under Contract No. DE-AC05-00OR22725. PLD, STEM, tr-KPFM and APT were conducted at the Center for Nanophase Materials Sciences, which is a U.S. Department of Energy Office of Science User Facility. N.B.G. gratefully acknowledge support from the U.S. National Science Foundation through grant DMR-1255379. This manuscript has been authored by UT-Battelle, LLC, under contract DEAC05-00OR22725 with the U.S. Department of Energy (DOE). The U.S. government retains and the publisher, by accepting the article for publication, acknowledges that the U.S. government retains a nonexclusive, paid-up, irrevocable, worldwide license to publish or reproduce the published form of this manuscript, or allow others to do so, for U.S. government purposes. DOE will provide public access to these results of federally sponsored research in accordance with the DOE Public Access Plan (http://energy.gov/downloads/ doe-public-access-plan)

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.8b00502. Details of the crystal structure analysis, transport analysis via tr-KPFM, yttrium distribution analysis from APT measurements, neutron scattering analysis, STEM imaging and analysis and DFT-modeling including ab initio molecular dynamics (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*(R.R.U.) E-mail: [email protected]. *(P.G.) E-mail: [email protected].

REFERENCES

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ORCID

Jilai Ding: 0000-0003-3905-8181 Janakiraman Balachandran: 0000-0002-2279-4441 Xiahan Sang: 0000-0002-2861-6814 Wei Guo: 0000-0002-9534-1902 Gabriel M. Veith: 0000-0002-5186-4461 Craig A. Bridges: 0000-0002-3543-463X Jonathan D. Poplawsky: 0000-0002-4272-7043 Raymond R. Unocic: 0000-0002-1777-8228 P. Ganesh: 0000-0002-7170-2902 Author Contributions ∇

Joint first author contribution

Author Contributions

P.G., J.B., and J.S.A. performed theoretical calculations, Y.Q. performed neutron measurements, R.R.U. and X.S. performed STEM measurements and data analysis, W.G. and J.D.P performed APT measurements and analysis, G.M.V, C.A.B. and C.M.R. prepared the sample targets and thin-films, J.D. was involved in all steps of sample preparation and characterization under the guidance of R.R.U and N.B.. P.G. oversaw the entire project. All authors, including N.B. were F

DOI: 10.1021/acs.chemmater.8b00502 Chem. Mater. XXXX, XXX, XXX−XXX

Article

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DOI: 10.1021/acs.chemmater.8b00502 Chem. Mater. XXXX, XXX, XXX−XXX