Interface Effects on the Ion Transport of Epitaxial Y2Zr2O7 Films

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Interface effects on the ion transport of epitaxial YZrO films Elisa Gilardi, Giuliano Gregori, Yi Wang, Wilfried Sigle, Peter A. van Aken, and Joachim Maier ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b00773 • Publication Date (Web): 25 Jul 2017 Downloaded from http://pubs.acs.org on July 26, 2017

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ACS Applied Materials & Interfaces

Interface effects on the ion transport of epitaxial Y2Zr2O7 films Elisa Gilardi †*, Giuliano Gregori†, Yi Wang†, Wilfried Sigle†, Peter A. van Aken†, Joachim Maier†

Max Planck Institute for Solid State Research, Heisenbergstr. 1, D-70569 Stuttgart, Germany KEYWORDS: Y2Zr2O7, ionic conductivity, interface effect, epitaxial films, mobility

Abstract

The systematic study of the ionic transport properties of epitaxial Y2Zr2O7 films with defective fluorite structure reveals an enhanced oxygen vacancy conductivity at the interface between the films and the MgO (110) substrate, which is characterized by a high density of misfit dislocations. This beneficial effect is discussed in terms of space charge and mobility effects.

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Introduction

Oxides with the chemical formula A2B2O7 (pyrochlores and defective fluorites) are well known ionic and mixed conductors and have been studied for many applications, such as catalysis, 1 materials for nuclear waste disposal, 2 thermal barrier coatings, 3 as well as potential candidates for electrolytes in solid oxide fuel cell thanks to their ability to conduct oxygen vacancies. 4 Pyrochlores and defective fluorites (which are sometimes described as [A,B]4O7) differ from each other by the order of the cationic sublattice and the occupancy of the anionic site. According to the Wyckoff notation, by taking the B site as the origin of a reference system, in the pyrochlore structure, the A cations sit on the 16c sites and the B ones occupy the 16d positions, oxygens sit on the 48f positions, while the oxygen vacant sites are located in the 8a positions.5 In the cation sublattice of a defective fluorite instead, all the cationic positions are crystallographically equivalent and can be occupied either by the A or the B cation. From this, it follows that the oxygen empty sites can sit in any anionic positions. In practice, a defective fluorite can be therefore treated as a heavily acceptor-doped fluorite, in which the concentrations of aliovalent cation and oxygen vacancy exceed the limit of a diluted situation. In such a case, the large concentration of point defects makes defect chemistry rather complex. In particular, the usually reasonable assumption according to which the mobility of the migrating point defects is independent of their concentration is not valid anymore: In such a situation the high concentration of point defects and their mobility are tightly interwoven to each other. It is worth noting that compared to an acceptor-doped fluorite structure (e.g. A-doped MO2, in which A M/ is negatively charged), in pyrochlores, these structural vacant oxygen sites are


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effectively neutral. As far as defect thermodynamics is concerned, they are to be considered as interstitial sites and are not involved in the oxygen vacancy migration process. The ratio between ionic radius of the cations r(A)/r(B) seems to be critical for determining the crystallographic structure of these oxides: for r(A)/r(B) > 1.26 the pyrochlore structure is more stable, while for r(A)/r(B) < 1.26 the defective fluorite structure prevails.6 These features strongly affect the electrical properties, since they are of importance for the oxygen migration mechanism.7 Oxygen migration in A2B2O7 compounds (pyrochlore and defective fluorite) proceeds via oxygen vacancy mechanism. The local atomic order obviously affects the migration energy barrier required to allow oxygen ions (or vacancies) to jump onto the next neighboring available crystallographic site.8 The pyrochlore phase usually exhibits lower migration energy and typically higher conductivity compared to the defective fluorite phase: for instance the difference in activation energy between the pyrochlore phase Nd2Zr2O7 (0.68 eV) and the defective fluorite phase Yb2Zr2O7 (1.44 eV) is quite remarkable.9 The electrical conduction properties of polycrystalline A2B2O7 ceramics have been extensively studied, particularly for the Gd2B2O7 system (B = Zr, Ti, Sn, Sm),

10-12

while for the defective

fluorite Y2Zr2O7 the literature is rather poor. Moon et al.13 investigated the Y2TixZr2-xO7 system and observed a general improvement of the ionic conductivity with increasing Zr/Ti ratio accompanied by a decrease of the migration energy. For the present study, the defective fluorite Y2Zr2O7 has been chosen as model system, because of the defined structure and the high chemical stability, as both cations are not sensitive to reduction and no chemical reaction with the substrate occurs during deposition or the conductivity measurement. Y2Zr2O7 can be treated as a highly Y2O3-doped ZrO2,



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2Y2O3 +2ZrZr× +O×O → 2YZr/ +VO•• +Y2Zr2O7 ,

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(1)

in which the concentration of oxygen vacancies is set by the electroneutrality condition

[YZr/ ] = 2[VO•• ] over a broad range of oxygen partial pressures and temperatures. It is noteworthy that for such a large dopant concentration, a strong defect association is expected, which can be expressed, e.g. by the following association reactions

 → (VO•• -YZr/ )• VO•• +YZr/ ← 

(2a)

 → (YZr/ -VO•• -YZr/ )X . VO•• +2YZr/ ← 

(2b)

Such reactions obviously affect the ionic transport properties (namely total conductivity and activation energy) of this compound. Thin films have been proven to be reliable systems for the characterization of the bulk electrical properties of the materials as well as possible interface effects,14-16 which can deeply modify the conductive properties at the interface between different layers. Various phenomena can alter the bulk properties of thin films, as for example strain effect, presence of dislocations or space charge effect. In this work, we investigate the electrical properties of Y2Zr2O7 epitaxial thin films of different thickness grown via pulsed laser deposition (PLD) on MgO (110) substrates. Quite remarkably, a clear positive interface effect is observed and it is correlated to the presence of a high density of misfit dislocations, which are required to accommodate the large lattice mismatch between film and substrate.


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The role of dislocations on the transport properties of oxides has recently gained increasing attention with controversial outcomes. While dislocations in YSZ single crystals seem not to affect the ionic transport,17 ‘superionic conductivity’ (~ 1 Ω-1 cm-1) has been reported in YSZ epitaxial single layers grown via pulsed laser deposition on MgO (110) and has been attributed to the high density of misfit dislocations under the assumption that each dislocation line can act as a fast transport pathway.18 On the other hand, dislocations can have the opposite effect, releasing the tensile strain deriving from the mismatch and lowering its beneficial effect on the oxygen transport as observed in multilayers of YSZ/Gd2Zr2O7 compared to multilayers of YSZ/CeO2.19 In this context, TEM investigations have shown that the interface between 8 mol.% yttria doped ZrO2 (8YSZ) single crystals and CeO2 epitaxial films exhibited a regular pattern of misfit dislocations and locally a strong reduction of cerium.20 Furthermore, computational studies on the role of dislocations in CeO2 21 and SrTiO3 22 highlighted that oxygen vacancies segregating in the core of edge dislocations are essentially immobilized therein. This is particularly obvious for dislocations constituting low angle grain boundaries in SrTiO3.23-25 Finally, it is worth noting that space-charge effects related to defect segregation in the dislocation cores, can substantially modify the electrical properties of oxides as it was observed in the case of TiO2 single crystals after having been plastically deformed at high temperatures. 26

Experimental section

Y2Zr2O7 oxide powder was prepared starting from ZrO(NO3)2 hydrate (Sigma Aldrich 99.99%) and Y(NO3)3 hydrate (Alpha Aesar 99.99%). The exact amount of water in the powder was determined by thermogravimetric analysis. The final oxide powder was obtained by heating and



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combustion of the aqueous solution of the two nitrates and 5% v/v of glycerol and calcination at 800 ºC for 5 hours. Chemical analyses on the powder have been performed by inductively coupled plasma optical emission spectrometry (ICP-OES, Ciros Spectro Analytical Instruments GmbH). The target for the pulsed laser deposition was pressed uniaxially (500 MPa) and sintered at 1700 ºC for 10 hours in air. The ICP-OES chemical analysis on the powders used for the target preparation indicated a negligible impurity content compared to the perceptible deviation of the yttrium-tozirconium atomic ratio from unity (Y/Zr = 0.97±0.02). Thin films were grown by PLD technique, using an excimer laser with wavelength of 248 nm (Coherent GmbH, Germany) on MgO (110) substrate of size 10 × 10 × 0.5 mm3 (Crystec GmbH). The energy density on the target was 1.5 J/cm2, with a pulse frequency of 5 Hz. The distance between target and substrate was 43 mm. During the film deposition the oxygen atmosphere was kept at 0.001 mbar and the temperature at 660 ºC (measured on the surface of the substrate with a pyrometer Heimann optoelectronics model KT19.99). After deposition the samples were annealed in 1 bar oxygen for 30 minutes at the same temperature as during deposition. The thickness of the films was measured using a scanning electron microscope Zeiss Crossbeam 1540 ESB. For the X-ray diffraction (XRD) analysis a Bruker D8 Diffractometer (Cu-Kα 1.5418Å) was used. The texture of the thin films was investigated using XRD pole figure analysis performed with a Philips X’Pert XRD – Co Kα radiation diffractometer with a wavelength of 1.78897 Å. The ϕ and ψ scan were recorded for 2 defined 2θ angles, namely at 40.50º and 69.0º. The software Carine 3.1 was employed to predict the position (ψ and ϕ degree) of the peaks of the selected plane.


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Transmission electron microscopy (TEM) analysis was carried out in order to investigate the microstructure of cross-sections of the films. The TEM specimens were prepared by tripod mechanical polishing down to 10 µm followed by argon ion beam milling in a stage cooled with liquid nitrogen. The transmission electron micrographs were acquired using a JEOL ARM 200CF Microscope equipped with a cold field emission electron source, DCOR probe corrector (CEOS GmbH), a 100 mm2 JEOL Centurio EDS detector, and a Gatan GIF quantum ERS electron energy-loss spectrometer. The microscope was operated at 200 kV, a semi-convergence angle of 20.4 mrad was employed giving rise to a probe size of 0.8 Å (1 Å for the analytical analyses), while 90–370 mrad collection angles was used to obtain high-angle annular dark-field (HAADF) images. For energy-dispersive X-ray spectroscopy (EDXS) measurements, Y-K and Zr-K lines were considered to quantify the concentration of the cations. The O-K-edge fine structure analysis was performed with a dispersion of 0.1 eV/pixel and with exposure time of 1 s/pixel. A collection semi angle of 56 mrad was used to record the EELS spectra. EELS and EDXS spectra were acquired simultaneously. After sputtering of two platinum electrodes (about 400 nm thick and 4.5 x 10 mm2 large, 1 mm distance between the electrodes) on the top of the films, their electrical properties were measured using an impedance spectrometer (Novocontrol Alpha-A High Performance Frequency Analyzer) operated at frequencies ranging between 2 MHz and 0.02 Hz by applying an A.C. voltage of 0.1 V. Impedance spectra were collected under oxidizing atmosphere and the pO2 was monitored by a commercial oxygen sensor (Rapidox 2100, Cambridge Sensotec). Typically, while bringing the samples from room temperature up to 700 °C, the samples were equilibrated in oxygen at 500 ºC for 9 hours (the equilibration process was monitored by recording impedance spectra every hour). Then, the temperature was increased with steps of 50 K and



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equilibrated for 4 hours at each temperature. Once the maximum temperature was reached (700 °C) the samples were kept 8 hours at 700 ºC. For the measurement of the activation energy the spectra were recorded by cooling the samples with temperature steps of 30 K with an equilibration time of 3 hours at each temperature. The impedance spectra were analyzed with the software ZView2 (Scribner Associates, Inc.). Hereafter the samples characterized after deposition are referred to as ‘pristine’ or ‘as deposited’ samples, while samples characterized after impedance measurement are referred to as ‘annealed samples’.

Result and discussion

Structural characterization

Thin films of various thicknesses were deposited on MgO (110) substrates exhibiting excellent epitaxy along the orientation (220) (Figure 1a). In the {001} pole figure the two sharp peaks at ψ=45º, confirm the epitaxial growth of the samples (Figure 1b). From the XRD pattern the outof-plane lattice parameter of 5.23 Å was calculated, which indicates an expansion of the lattice in the direction perpendicular to the surface of about 0.6% (compared to the value of 5.21 Å reported in the literature).27 AFM analysis (not shown here) revealed the surface of these films to be smooth with an average root mean square roughness of 0.5 nm (which corresponds roughly to one unit cell).


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Figure 1 Structural characterization of a Y2Zr2O7 thin film grown on MgO (110). a) XRD pattern b) {001} pole figure (2θ = 40.5°)

TEM analysis

TEM analyses performed on cross-sections of the Y2Zr2O7 films confirmed the high degree epitaxy of the films (Figure 2a) in spite of the large lattice constant mismatch (–23%) between the substrate (aMgO = 4.21 Å) and the film (aYZO = 5.20 Å). TEM micrographs acquired from an as-prepared sample (Figure 2b) indicate the presence of a high density of misfit dislocations at the film/substrate interface, which are structurally required to accommodate the different lattices. Interestingly, the exposure to high temperature during the conductivity measurements (700 °C for 8 hours) induces a structural rearrangement of the interface. While in the pristine sample the interface is flat and the average distance between neighboring misfit dislocations is 4.5 atomic planes, after the heat-treatment, the interface exhibits a saw-tooth-like profile and the average distance between adjacent misfit dislocations is increased to 5.4 atomic planes (Figure 2b and c). Such a change can be rationalized by considering theoretical calculations on MgO (110), according to which a faceted MgO surface is expected to be more energetically stable than the flat one.28 9 ACS Paragon Plus Environment


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Figure 2. a) TEM micrograph of a thin film of Y2Zr2O7 on MgO (110). The inset shows an electron diffraction pattern of a selected area including both substrate and film, showing the epitaxial character of the sample. b) Top: HAADF image of the interface, Bottom: Fourierfiltered image clearly showing the misfit dislocations created at the interface in order to release the stress deriving from the large mismatch between substrate and film (–23%). The average number of atomic planes between two subsequent misfit dislocations is indicated. The inset shows a magnified detail of the bottom image highlighting two neighboring dislocations at the interface. (c) Top: HAADF image of the interface after 700 °C annealing revealing the saw-


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tooth-like shape of the interface. Bottom: Fourier-filtered image of the sample showing a decreased number of misfit dislocations.

EDXS analysis

EDXS analysis was performed on two samples, one directly after deposition (Figure 3a,c) and one after exposure to 700 ºC, as described previously (Figure 3 b,d). Particular attention was paid on the yttrium-to-zirconium ratio, which defines the concentration of oxygen vacancies in the material. As shown in Figure 3 a,b, the Y3+/Zr4+ ratio is constant throughout the sample and a closer inspection of the interface (Figure 3 c,d) indicates that also in the first few nanometers near the substrate/film boundary the ratio remains unchanged. An overlap of the signal of Mg2+ from the substrate and the signals of Y3+ and Zr4+ from the film can be noticed in the first half nanometer near the interface. At first, this could suggest a certain cationic intermixing occurring between film and the substrate. However, in the case of Mg diffusing into the zirconate layer one would expect to detect a deviation of the Y-to-Zr ratio close to the interface due to Mg preferentially substituting Y or Zr (an analogous reasoning can be applied for Y and Zr diffusing into MgO). Since this is not observed and since the width of the intermixing region is comparable to the roughness of the substrate, we conclude that this overlap most likely stems from the finite roughness of the substrate rather than an evidence of a substantial interdiffusion. The decrease of the EDXS net counts close to the surface of the sample arises from the thickness variation of the TEM specimen (along the electron beam direction). The cross-sectional TEM specimens were wedge polished with an angle of about 1.5 degree, and then ion milled from one side. Thus, the final TEM specimens are wedge shaped yielding to a cross-section, which is



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thinner in proximity of the surface than in portion close to the substrate. The measurement of the as-deposited sample (Figure 3 a) was performed in a relative thick area, which makes the influence of specimen thickness variation less pronounced.

Figure 3. EDX analysis of the overall sample normal to the interface of a) the pristine sample and b) the sample after annealing at 700 °C. The ratio between Y and Zr is constant throughout the complete sample. c) linescan of the first nanometers near the interface in steps of 1 Å of the as deposited sample and d) after annealing. Y/Zr ratio is constant also near the interface. To be noticed the overlap of Mg with Zr and Y in the first nanometer. Counts are normalized with respect to the maximum in each linescan.

Electron energy loss spectroscopy (EELS)


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The results of the EELS analysis of the oxygen K-edge are summarized in Figure 4. A series of five EELS spectra were acquired perpendicularly to the interface starting with the MgO substrate (position 1 in Figure 4a) and ending with the Y2Zr2O7 film (position 5). As shown by the EELS spectra reported in Fig. 4b, it is evident that the fine structure of these O-K-edges changes perceptibly along the acquisition direction. After removal of the background contribution, the intensity of the spectra across the interface I O -K ( x ) was fitted using a multiple linear least-squares (MLLS) method, which is based on the

linear combination of the intensities of the characteristic edges resulting from MgO and Y2Zr2O7 (spectra 1 and 5, respectively)

Y 2 Zr2 O 7 , I O -K ( x ) = a ( x ) ⋅ I OM-gO K + b ( x ) ⋅ I O -K

(3)

where a and b are numerical coefficients (depending on the position x) ranging between 0 and 1. It is worth noticing that the O-K-edge spectrum at the interface region (spectrum 3 in Figure 4b) of the annealed sample is perfectly fitted by the linear combination of the characteristic spectra of MgO bulk and Y2Zr2O7 bulk. The MLLS fit of the spectrum acquired from the interface of the pristine sample leads to a remarkable residual signal at the interface as shown in Figure 4c indicating the presence of additional O-K-edge spectral features. In order to further analyze this finding, the spectral range 528–542 eV was decomposed into three Gaussians by using nonlinear least square (NLLS) fitting for all the spectra acquired along the line-scan (Figure 5a). The variation of the intensity of the first Gaussian (centered at 530 eV) as a function of the distance from the interface is shown in Figure 5b and c in red. In order to compare them to the profile of the O-K-edge spectra of the Y2Zr2O7 environment (obtained



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previously by fitting the all spectra to spectra 1 and 5 along the line scan, blue curves in Figure 5b and c), the intensity profile of the first Gaussian peak was scaled to 1 in the region of bulk Y2Zr2O7. In the pristine samples the two profiles have a remarkable difference within the first nanometer close the interface (Figure 5b) while for the annealed sample the two intensity profiles match perfectly (Figure 5c). From this, it is clear that, within a layer of approximately 1 nm (note that the ability of resolving the width of this layer depends on the sensitivity of the EELS technique), oxygen atoms at the interface of the pristine samples are surrounded by a different chemical environment compared to the bulk situation. As stated above, this result cannot be explained by a linear combination of the signals stemming from MgO and Y2Zr2O7, respectively nor, as shown by EDXS, to a deviation of the Y/Zr ratio from the bulk value. Therefore, these unique features of the O-K-edge are suggestive of a very specific oxygen stoichiometry at the interface, which we assign to the presence of oxygen vacancies (cf. the results of the impedance spectroscopy measurements discussed below).


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Figure 4. (a) Cross-sectional STEM-HAADF micrograph showing the positions (numbered open circles) where the EELS spectra (b) were acquired from the pristine sample. Position 1 corresponds to the bulk MgO substrate while position 5 refers to the bulk Y2Zr2O7 film. (c) Residual signal resulting from the subtraction of the MLLS fit of bulk MgO and bulk Y2Zr2O7 spectra from the spectra acquired at the interface (position 3). The black line refers to the annealed sample and the red line to the as-deposited one.



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Figure 5. (a) decomposition of the O-K-edge by means of 3 gaussians. O-K edge intensity as a function of the distance from the interface for (b) the pristine sample and (c) the annealed one. The red line indicates the intensity of the first O-K pre-edge peak at 530 eV resulting from the NLLS fitting; the blue line represents the integrated intensity of the whole O-K edge spectrum at each position.

Impedance spectra

The impedance spectra for determining the electrical conductivity and the corresponding activation energy were recorded after equilibration for 9 hours at 500 ºC and 8 hours at 700 ºC of 16 ACS Paragon Plus Environment

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the samples at constant oxygen partial pressure. The spectra were fitted using an equivalent circuit consisting of an electrical resistance in parallel with a constant phase element (Q) as shown in Figure 6a. The analysis, repeated in different gas environments, confirmed the independent behavior of the electrical conductivity from the oxygen partial pressure, indicating oxygen vacancies as the majority charge carriers (Figure 6b). The dependence of the conductivity on the temperature of a thin film 128 nm thickness is shown in Figure 6c. Conductivity and activation energy values (1.44 eV ± 0.02 eV) are in agreement with data reported in the literature for ceramic pellets. For example Yamamura et al.9 reported a total conductivity of 1.8 10-4 S cm-1 at 800 °C and an activation energy of 1.3 eV.



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Figure 6. (a) Impedance spectra (Nyquist plot) of Y2Zr2O7 thin films on MgO (110) of different thickness at 500°C (b) Dependence of the ionic conductivity on the oxygen partial pressure of a 300 nm thick film at 500°C. The experimental error of the conductivity data derives from the sum of the error of the fit of the experimental spectra and the uncertainty stemming from the 18 ACS Paragon Plus Environment

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determination of the electrodes geometry (electrodes length and their distance from each other). c) Temperature dependence of the conductivity of a Y2Zr2O7 thin film with 100 nm thickness.

Thickness dependence of the conductance

From the study of the electrical properties of thin films, in particular the epitaxial ones, which are free from any grain-boundary contribution, it is possible to obtain information about film/substrate interfacial effects. For this purpose, it is convenient to normalize the measured conductance G by the geometrical dimensions perpendicular to the growth directions l1 and l2 according to29

Y // =G //×

l1 l2

.

(4)

It is important to note that by doing so, Y can be then expressed by two terms: the bulk contribution Y∞ and the film-substrate interface contribution ∆Y//:

Y// =Y∞ +∆Y// =σ∞L+∆Y//

(5)

with L being the sample thickness. This means that in the Y// vs L diagram shown in Figure the slope represents the bulk conductivity σ∞, while the extrapolated intercept on the ordinate axis gives the value of ∆Y//. As the charge of the majority mobile defect is positive, ∆Y// > 0 means an enrichment of oxygen vacancies or an improvement of their mobility, whilst ∆Y// < 0 indicates a local depression of the ionic transport.



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It is noteworthy that these samples show a clear interface contribution to the conductance, resulting from the positive intercept ∆Y// (2.94×10-11 ± 1.9×10-11 S, at 500 °C). Such ∆Y// is clearly larger than the normalized conductance resulting from the measurements of the bare substrate (grey area) and thus indicates the occurrence of an interface effect between film and substrate. Interestingly, the positive intercept could be observed at all the temperatures considered between 500 ºC and 700 ºC. After exposure to the maximum temperature (700 °C) for 8 hours, the bulk conductivity very slightly decreases (–20 %), while ∆Y// decreases more pronouncedly (–35 %) (See Figure and Table 1).

Figure 7. Normalized conductance Y// vs thickness of epitaxial Y2Zr2O7 thin films grown on MgO (110) at 500 ºC, during a heating (blue circles) and a cooling cycle (red circles). The grey area corresponds to the normalized conductance of the bare substrate. In the inset: magnification of the intercept on the y axis, which corresponds to the value of ∆Y//. The error bar on the y axis has been obtained taking into account the error resulting from fitting the impedance spectra and the error deriving from the geometry of the electrodes. The error bar on the x axis results from the thickness measurements performed with the SEM. For the intercept on the y axis confidence intervals (within the 90% of probability) are reported.


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Table 1 Activation energy (in eV) of the interface and bulk contributions calculated from the plot of the sheet conductance versus the thickness. For comparison, the conductivity activation energies of the total film contributions (of different thicknesses) calculated on the same range of temperature are also included. Confidence intervals (within the 90% of probability) are reported, standard errors have been obtained taking into account the error resulting from fitting and the error deriving from the geometry of the electrodes.

Interface

Bulk

Film 70 nm

128 nm

328 nm

425 nm

Before annealing

1.5±0.3

1.5±0.2

1.5±0.3

1.4±0.3

1.5±0.5

1.5±0.5

After annealing

1.5±0.3

1.5±0.1

1.5±0.4

1.5±0.4

1.5±0.5

1.5±0.4

The variation of σ∞ (from 2.5×10-6 S cm-1 before exposure to 700°C to 2×10-6 S cm-1 after annealing) is probably due to an artifact, namely the dewetting of the Pt electrodes resulting in a variation of the electrode geometry with time. 30 Obviously these points would also affect the value of ∆Y//, however as the variation of ∆Y// is more pronounced than the change of the bulk conductivity, a substantial modification of the interface properties occurred during/after exposure to 700 °C. The positive enhancement of the conductance (∆Y//) is naturally linked to the specific microstructure of the MgO/Y2Zr2O7 interface, which exhibits a high density of misfit dislocations. It was recently demonstrated by theoretical studies as well as oxygen diffusion experiments on related oxygen vacancy conducting materials (namely CeO2 and SrTiO3) that



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dislocation cores tend to hinder the oxygen ion transport, making the ionic conduction through these pathways very improbable.21-22 It seems therefore reasonable to rule out a possible preferential migration along the dislocation cores also for the present case. Notably, such a large concentration of extended defects makes the interface very likely to be charged. The segregation of positively charged oxygen vacancies in the dislocation core would deplete the oxygen vacancies outside the dislocations, because of the excess of positive charge thus hindering the ionic transport locally. Conversely, a negatively charged interface would instead lead to a local enhancement of the mobile positive charge carriers, which is consistent with the experimentally observed positive intercept ∆Y//. In this context, it is noteworthy that theoretical calculations on Schottky defects in MgO estimated the excess energies of magnesium vacancies and oxygen vacancies at the edge // dislocations to be -1.7 eV and -1.5 eV, respectively suggesting preferential segregation of VMg

••

into the dislocation core (rather than of VO ).31 Thus, the accumulation of negatively charged magnesium vacancies in the dislocation core would result in an excess of negative charge at the interface, and hence to an enrichment of mobile oxygen vacancies within the adjacent spacecharge region. The effect of different doping content on the transport properties of YSZ has been studied in the past years, in particular for doping level up to 30%

32-33

. However, due to the high doping

content in Y2Zr2O7 the general assumption that free carrier concentration and mobility are independent is not valid anymore (cf. for example ref. D.Y. Wang et al, Solid State Ionics 2, 95 (1981)),34 which makes it more difficult to predict the electrical properties of the material in case of oxygen vacancies enrichment or depletion at the interface with the MgO substrate.


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According to the space-charge effects

35

the accumulation of oxygen vacancies in the space

large layer corresponds to:

[VO•• ]0 / [VO•• ]∞ = exp−2 e∆φ0 /( kBT ) ,

where

(6)

∆φ0 is the space-charge potential and kBT the Boltzmann term. ••

/

Considering the already large concentration of oxygen vacancies ( [VO ]∞ = [YZr ]∞ / 2 ), the local •• enhancement of [VO ] is rather limited: As the number of available oxygen lattice sites per unit

cells is 8, the ratio between the concentration of vacancies at the interface [VO•• ]0 and their bulk concentration [VO•• ]∞ cannot exceed 8. •• This means that if [VO ]0 is eightfold larger than the bulk value, ∆φ0 is rather small and on the

order of –70 mV at 500 °C (we note that this represents the lower boundary for the space-charge potential).

Allowing also for a varied mobility of the charge carriers and approximating the space charge •• zone by a rectangular profile of concentration [VO ]0 and extent l , we may write for the excess

conductance:

∆Y // ≅ 2eu0{[VO•• ]0 − [VO•• ]∞}l .

(7)



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(Note that for the determination of ∆Y// the conductance of the substrate Y MgO = 4.2×10-12 S is taken into account). By comparison with the O-K-EELS analysis, we can estimate the width the space charge zone to be 1 nm. Even in this limit case the accumulation of oxygen vacancies alone is not sufficient to explain the enhanced ionic transport at the interface; rather an improved mobility is required. For •• •• example, if [VO ]0 = 8[VO ]∞ then u0 ≅ 15 u∞ results. •• As the above expression clearly overestimates the total [VO ]0 content, the result for u 0 is the

lower bound. If we however more realistically assume strong association leading to a much lower free carrier concentration (with a correspondingly much higher mobility), space charge effects might explain the conductance increase. In this context, it is worth noting that this interdependence between concentration and mobility is however inherent to such highly nonideal situations and it is not advisable to separate them in the usual way. A clear information on the degree of space charge effects would require a measurement of the core charge which we do not have. We note that the space-charge effect together with the improved mobility at the interface should correspond to different activation energy of the interfacial conductance compared to the bulk ionic transport. However, experimentally we do not see any perceptible difference of these activation energies. This can be attributed to the rather large error (±0.3 eV) associated with the activation energy of the interfacial conductance that can overshadow the actual variation of the activation energy. After annealing, the lower interfacial conductance is suggestive of a less pronounced effect at the substrate/film boundary: It is however again difficult to establish whether this is mostly due to

a

mobility

change

or

to

a

modification

of

the

space-charge

situation. 24

ACS Paragon Plus Environment

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Conclusion

Y2Zr2O7 epitaxial thin films were successfully grown on MgO (110) substrate. From the study of samples with different thicknesses a positive interface effect between the films and the substrate appeared at different temperatures and it decreases upon exposure to 700 ºC. TEM analyses reveal a high number of misfit dislocations, which are naturally present at the interface, in order to accommodate the strain deriving from the large lattice mismatch between film and substrates. EELS analysis of the O-K-edge pointed out a different chemical environment at the film/substrate interface in the pristine sample compared to the annealed one. Notably, the comparison of the EELS results with the data obtained from impedance spectroscopy analysis suggests the oxygen vacancies between film and substrate, which owing to the strongly non-ideal defect chemistry also involves mobility effects. If more realistically one takes into consideration

bulk defects association leading to a lower bulk free carrier

concentration (with a correspondingly higher bulk mobility), then space charge effects might explain the conductance increase. This interdependence between concentration and mobility is however inherent to such highly non-ideal situations and it is not advisable to separate them in the usual way.

AUTHOR INFORMATION Corresponding Author *[email protected] Present Addresses



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†Laboratory for Multiscale Materials Experiments, Paul Scherrer Institut, 5232 Villigen, Switzerland Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Founding Sources Y.W., W.S. and P.A.v.A. wish to acknowledge the European Union Seventh Framework Program [FP/2007/2013] for funding under grant agreement no 312483 (ESTEEM2). ACKNOWLEDGMENT The authors wish to thank G. Cristiani and B. Stuhlhofer for their technical support for pulsed laser deposition. U. Salzberger is thanked for the TEM samples preparation, M. Dudek is thanked for the pole figures acquisition. B. Fenk is thanked for the SEM analysis. The Central Scientific Facility for chemical synthesis of the Max Planck Institute for Intelligent System Stuttgart is thanked for the chemical analysis. R. Merkle is thanked for helpful discussions.


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Table of contents


Interface Effects on the Ion Transport of Epitaxial Y2Zr2O7 Films

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