Nanoscaffold WO3 by Kinetically Controlled Polymorphism - Crystal

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Nanoscaffold WO by kinetically-controlled polymorphism Jaeseoung Park, Hyojin Yoon, Si-Young Choi, and Junwoo Son Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.8b01551 • Publication Date (Web): 06 Dec 2018 Downloaded from http://pubs.acs.org on December 6, 2018

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Crystal Growth & Design

Nanoscaffold WO3 by kinetically-controlled polymorphism

Jaeseoung Park1, HyojinYoon1, Si-Young Choi1,a) and Junwoo Son1,b)

1 Department

of Materials Science and Engineering, Pohang University of Science and Technology (POSTECH), Pohang 37673, Republic of Korea

0

ACS Paragon Plus Environment

Crystal Growth & Design 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Abstract Several metastable polymorphs, diverse crystal structures with a same chemical formula, could be accessed through epitaxial interfaces with low energy between epitaxial film and symmetrymatched substrates. Here, we fabricated tungsten oxide (WO3) nanoscaffolds composed of hexagonal WO3 (h-WO3) / monoclinic WO3 (m-WO3) polymorphs, and modulated the proportions of these phases by tuning W arrival rate during epitaxial growth. The WO3 nanoscaffold is vertically aligned with coherent interphase boundaries between h-WO3 and m-WO3; this structure leads to the persistence of h-WO3 despite its meta-stable characteristic. The coherent interphase boundaries give rise to the unexpected six-fold m-WO3 phases and a distinct lattice response with intercalated hydrogen defects. This WO3 nanoscaffold with different fractions of two polymorphs would be a good model system to systematically study intercalation chemistry and its application with WO3 polymorphs.

a)

[email protected]

b)

[email protected] 1


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Introduction Due to various possibilities to arrange polyhedra in transition metal oxides, they can exhibit numerous crystal structures (polymorphs) that have the same chemical formula. Although polymorphs with distinct crystal structure have been attracting the considerable attention due to their unique properties and practical applications1-6, many are thermodynamically metastable in ambient conditions, and cannot even exist in bulk form. Non-equilibrium epitaxial growth technique has yielded thin films of these metastable polymorphs by epitaxial stabilization on symmetry-matched substrates with precise control of growth condition2, 7, 8. In sufficiently-thin films, similar interfacial registry of metastable polymorph films with appropriate substrates provides a sufficiently low interfacial free energy EI that can overcome the cost of volume free energy EV between polymorphs. As examples, anatase-TiO2 phase that was epitaxially stabilized on (001) perovskite surfaces, showed better photocatalytic effects than thermodynamically stable rutile-TiO23; and metastable bronze-phase VO2(B) stabilized on (001) SrTiO3 showed potential as an energy-storage material due to its open atomic framework and good room-temperature conductivity2, 8. Tungsten oxide (WO3) has been widely studied due to its strong potential applications in energy-related devices, such as smart windows9-12, gas-sensors13, 14, and hydrogen production15, 16. WO3 also can be crystalized into distinct polymorphs. Monoclinic WO3 (m-WO3, Fig. 1b) with Asite vacant distorted perovskite structure is thermodynamically stable in ambient conditions. In contrast, Hexagonal phase (h-WO3, Fig. 1c) is metastable; it has a unique crystal structure that includes an empty intercrystalline tunnel (diameter ~ 5.36 Å) along the c-axis, and trigonal 2



nanochannels in a framework of corner-sharing WO6 octahedra; h-WO3 can be stabilized with intercalated ions or small molecules by anisotropic diffusion through this tunnel in bulk form17, 18. The high surface area of the tunnels facilitates electrochemical reactions and ionic transport in hWO3, so it achieves higher specific pseudocapacitance and more accessible electrochemical active sites than does m-WO34, 19. Motivated by this facilitated intercalation chemistry, h-WO3 nanoparticles show great potential for various applications, such as electrochromic windows20, cathodes for lithium batteries21, selective ion transfer and gas absorption4, mixed protonic-electronic conductor19, and even superconductivity after heavy alkali metal ion doping22. Despite the importance of diverse functionalities in h-WO3 polymorph, its growth kinetics of single crystal h-WO3 has not been systematically studied on a large scale, and its intercalation chemistry has not been studied due to absence of a strategy to synthesize single-crystals without intercalated ions. In this study, we systematically demonstrate an h-WO3 / m-WO3 nanoscaffold by balancing the epitaxial stabilization of h-WO3 on the lattice-matched substrates against the thermodynamic instability of metastable h-WO3 (Fig. 1a). The phase fractions of the two phases were kinetically controlled by changing the growth rates (laser frequency 0.5 ≤ f ≤ 10 Hz) at growth temperature of 500 °C, at which the metastable h-WO3 phase tends to transform to thermally-stable m-WO3 phase. This WO3 “nanoscaffold” consists of two WO3 polymorphs that are highly oriented along the out-of-plane direction as well as in-plane direction with an unexpected crystal symmetry due to the coherent interfaces between two polymorphs along the vertical directions. High-resolution X-ray diffraction and atomic-resolution electron microscopy reveal that the WO3 nanoscaffolds 3


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have vertical interphase boundaries between h-WO3 and m-WO3. Unexpected six-fold m-WO3 phases grow by forming coherent interphase boundaries with h-WO3. Unlike m-WO3 single phase, these WO3 nanoscaffolds showed irreversible lattice response by hydrogen intercalation, which occurs because strong interactions between hydrogens and h-WO3 lattice increase the energetic favorability of intercalated h-WO3 phase compared to pristine h-WO3 phase.

Methods The growth of h-WO3 / m-WO3 nanoscaffolds The 60-nm thick h-WO3 / m-WO3 nanoscaffolds were grown by pulsed laser deposition (PLD) on step-treated (111) YSZ substrates. For comparison, m-WO3 films with thickness of 60 nm were grown on (0112) Al2O3 substrates. A stoichiometric WO3 (99.9%, Sigma Aldrich) pellet was prepared as a PLD target by cold isostatic pressing with a pressure of 2000 kg/cm2 for 2 min, then sintering in air at 900 °C for 24 h. Before the growth of h-WO3 / m-WO3 nanoscaffolds, the (111) YSZ substrate was annealed in air at 1400 °C for 2 h to obtain and atomically flat step-terrace structure23. The samples were loaded into a PLD chamber, which was then evacuated to a base pressure of ~ 1 x 10-6 Torr; then the rotating WO3 target was ablated by focusing KrF excimer laser (λ = 248 nm) with a fluence of 1 J/cm2 at substrate temperature of 500 °C and oxygen pressure of 200 mTorr. The growth rate was controlled by changing repetition rates f of the laser shot from 0.5 Hz to 10 Hz, and the rate per single laser was accurately determined to be ~ 0.13 nm/shot using x-ray reflectivity measurement. 4



Materials characterization For structural characterization, high-resolution X-ray scattering measurements (symmetric 2θ- scans and skew-symmetric φ-scans) were conducted using X-ray diffraction (D8 discover, Bruker) and synchrotron radiation (λ = 0.123984 nm, energy = 10 keV) at the 3D XRS beamline of Pohang Light Source-II. For in-situ monitoring of out-of-plane lattice parameters during hydrogenation and dehydrogenation, a home-built in situ chamber was used. During hydrogenation and dehydrogenation, the forming gas (H2 (5%) / Ar (95%)) and dry air gas, respectively, was introduced into the chamber at flow rates of 500 sccm under 1 atm. The phonon modes of WO3 nanoscaffolds were obtained by confocal Raman spectroscopy (Alpha300 R, WITec) using a 532-nm laser. To achieve precise measurement, each spectrum result was determined as the average of ten independent measurements obtained during 1 min. The sample for STEM analysis was prepared using a Focused Ion Beam (FIB). The specimen was ion-milled using a 0.1 ~ 0.5 keV Ar beam (PIPS II, Gatan) to thin the sample and to minimize surface damage imparted by the FIB process. STEM analysis was performed using a 200-kV STEM (JEM-2100F, JEOL) equipped with an aberration corrector (CEOS GmbH). The optimal size of the electron probe was ~ 0.8 nm. The A 9C (23.2 pA) probe was used to acquire HAADF and ABF images. The collection semi-angles of the HAADF detector were adjusted from 70 to 250 mrad to get clear Z-sensitive images. ABF images were obtained using collection angles from 12 to 24 mrad; the raw images were processed using a band-pass Wiener filter with a local window to reduce background noise (HREM research Inc.). STEM image simulation was performed using QSTEM (free software developed by Christophe Koch, Humboldt University). 5


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Hydrogen spillover of WO3 nanoscaffolds with Pt nanoparticles Pt hydrogen spillover was used to efficiently inject hydrogens into the WO3 nanoscaffolds at low temperature24. Nano-sized Pt islands were deposited on the WO3 layers by rf magnetron sputtering, then the samples were annealed under forming gas atmosphere at 100 °C. This low-temperature hydrogen spillover method was intended to minimize the formation of oxygen deficiencies in the nanoscaffolds. For dehydrogenation, the hydrogenated samples were cooled to room temperature under forming-gas atmosphere, then blown using dry air in in situ measurement chamber.

Results and Discussion Epitaxial stabilization of h-WO3 / m-WO3 nanoscaffold Pulsed laser deposition (PLD) was used to epitaxially stabilize h-WO3 films on (111) YSZ substrate, because they both have hexagonal crystal symmetry, and because the lattice constant of (111) YSZ is close to twice (2 x 𝑑110 of YSZ ~ 7.268 Å) of that of h-WO3 (~ 7.298 Å) (Fig. 1c, d)7, 25: Symmetric 2θ-ω scans of high-resolution X-ray diffraction (HRXRD) showed only one Bragg’s peak in 60-nm thick WO3 films that had been synthesized at the highest growth rate (f = 10 Hz at 500 °C) (Fig. 2a). The Bragg’s peak occurred at 18.38°, which corresponds to interplanar distance of 7.762 Å; This value is similar to 2 x 𝑑0001 of h-WO3 thin films25 (7.76 Å) or powder17 (7.798 Å) (Fig. 2a). In addition, the clear Laue fringes and the single unit-cell step height were observed by HRXRD and by atomic force microscopy, respectively (Fig. S1, 2), which indicate sharp interfaces and atomically flat surfaces between the substrate and WO3 films with high quality. 6



To verify the epitaxial relationship between WO3 layer and YSZ substrates, skewsymmetric φ-scans of h-WO3 and YSZ planes were conducted using WO3 films grown at f = 10 Hz (Fig. S3). (Due to different crystal symmetries, so to avoid confusion, the sets of planes are denoted as (h k l)YSZ for YSZ, (h k i l)h for h-WO3 and (h k l)m for m-WO3.) The φ-scan spectrum on (202 2)h plane showed six-fold symmetry and matched well with that of the (220)YSZ plane of YSZ substrate; this similarity confirms that the (2022)h plane of h-WO3 phase was parallel to the threefold (220)YSZ plane of YSZ substrate (Fig. 2b). Indeed, the nearest inter-oxygen distance of outof-plane (0001)h-faceted h-WO3 (3.649 Å) is quite similar to that of (111)YSZ-faceted YSZ (3.634 Å); oxygen atoms except those in YSZ aligned with hexagonal tunnels in h-WO3 are shared perfectly at the interface between these two hexagonal materials (Fig. 1c, d); the epitaxial relation between h-WO3 phase and YSZ substrate was confirmed as (0001)h [1000]h h-WO3 || (111)YSZ [11 0]YSZ YSZ. These lattice-matched substrates decrease the EI between h-WO3 and YSZ substrates, so metastable h-WO3 is stabilized even without any intercalations; this lattice matching is the key to stabilizing the metastable hexagonal phase. In addition to six-fold peak of φ-scan in h-WO3, the φ-scan (Fig. 2b) of the (202)m plane of m-WO3 also shows tiny six-fold symmetry peaks, which indicate that h-WO3 and m-WO3 coexist in the films; the formation of the m-WO3 phase was inevitable even at high f despite the maximization of epitaxial stabilization of h-WO3 by the (111)YSZ YSZ substrates with the same symmetry. Interestingly, the six-fold symmetry of the (202)m m-WO3 peak is surprising given that the A-site vacant perovskite crystal structure of the (202)m plane of m-WO3 is supposed to have two-fold or four-fold symmetry (Fig. 2b, Fig. S3); the observation is consistent with the presence 7


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of three different sets of m-WO3 domains, each rotated 60° to the others. Each set of m-WO3 domains was rotated azimuthally by 30° from the (2022)h plane of h-WO3 and the (220)YSZ plane of YSZ. Together, these results indicate that the epitaxial relation among h-WO3, m-WO3, and YSZ is (0001)h [1000]h h-WO3 || (001)m [100]m m-WO3 || (111)YSZ [110]YSZ YSZ. m-WO3 and YSZ have completely different in-plane symmetry (Fig. 1b, d), so the aligned six-fold symmetry of m-WO3 by direct growth of (111)YSZ YSZ substrate was not expected; we hypothesize that sixfold m-WO3 phase was self-assembled and stabilized by the epitaxial relationship with h-WO3 phase on YSZ substrates.

Atomic scale analysis of h-WO3 / m-WO3 nanoscaffold with vertical interphase boundaries The atomic-scale crystal structures and the interphase boundaries of WO3 nanoscaffold were visualized by atomic-scale electron microscopy. HRTEM image and the selected area diffraction pattern (SADP) of the h-WO3 / m-WO3 nanoscaffold (f = 10 Hz) along the [110]YSZ zone axis (Fig. 3a) show the presence of both WO3 polymorph phases and YSZ substrate. The mWO3 phase shows a pseudocubic diffraction pattern, whereas the h-WO3 phase shows an alternating diffraction contrast and anisotropic crystal symmetry perpendicular to the [1000]h zone axis. The phase coexistence of WO3 polymorphs was visualized in magnified HAADF-STEM image along the [110]YSZ zone axis (Fig. 3b). Since both polymorphs include the same atomic mass Z of heavy cations (W), Z contrast between h-WO3 and m-WO3 is not prominent. However, the polymorphs show distinct sequences of atomic contrast in the [110]YSZ zone axis parallel to [1000]h || [100]m direction. Abrupt vertical interfaces separate two regions: (1) in which a similar 8



atomic contrast repeats, and (2) in which a strong and a weak atomic contrast alternate along the in-plane direction. The plane view of h-WO3 / m-WO3 nanoscaffolds (Fig. 3d) along the out-ofplane direction ([111]YSZ) suggests that (1) could be assigned to ReO3-type m-WO3, because it has the same atomic density along the [100]m or [010]m direction. In contrast, (2) could be assigned to h-WO3 phase, because this phase has an open tunnel along the [0001]h direction; this tunnel results in weak atomic density columns from the view of [1000]h directions (Fig. 3d). The h-WO3 phases were coherently clamped by the YSZ substrates, because the nearest inter-oxygen distance of the (0001)h h-WO3 plane (3.649 Å) is quite similar to that of (111)YSZ YSZ plane (3.634 Å). However, edge dislocations toward YSZ substrates exist under the m-WO3 phase because of the larger inter-oxygen distance in m-WO3 than in YSZ substrate (Fig. 1b, d): Due to crystallographic difference, the larger strain energy between m-WO3 phase and YSZ substrate should be released by forming the dislocations. To quantitatively analyze the atomic structures of the h-WO3 / m-WO3 nanoscaffold (f = 10 Hz), the magnified HAADF- and ABF-STEM analyses along [110]YSZ zone axis were examined closely (Fig. 3c). The alternating atomic column density and epitaxial relationship between h-WO3 and YSZ suggests that the [110]YSZ direction (zone axis of Fig. 3a-c) is parallel to [1000]h25. By dividing the W –W distance by the number of unit cells, 𝑑110 of h-WO3 was calculated as 6.318 Å from the HAADF image; this value is comparable to that of bulk h-WO3 (6.32 Å). As expected from the phase-resolved φ-scans of HRXRD (Fig. 2b), the boundaries between h-WO3 and m-WO3 were coherent without any dislocations; the perfect atomic arrangement in the vertical direction is ascribed into the small lattice mismatch between m-WO3 9


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(𝑑001 = 7.678 Å) and h-WO3 (2 × 𝑑0001

= 7.798 Å) along the out-of-plane direction ([0001]h ||

[001]m) (Fig. 1b, c). Due to small lattice mismatch between m-WO3 and h-WO3, they share oxygen atoms along the vertical direction and develop abrupt and coherent interphase boundaries. HAADF images suggest that the structure is an ordered intergrowth of metastable h-WO3 slabs joined to mWO3 slabs at coherent vertical interfaces.26 Because of lower EI at the coherent interface between two polymorphs (m-WO3, h-WO3) than between m-WO3 and substrates, the m-WO3 phase should nucleate preferably and grow laterally on the h-WO3 phase rather than on YSZ substrates. Therefore, epitaxial growth of h-WO3 on YSZ substrates should guide the formation of ordered h-WO3 / m-WO3 nanoscaffolds with vertical interfaces. In vertical heteroepitaxial nanocomposites, in which abrupt interfaces between component phases in nanocomposites exist along the vertical direction, efficient strain coupling and coherency of the interfaces occur between component phases, rather than coupling between phase and underlying substrates, so the film thicknesses and lattice constants of component phases affect the tuning of strain and physical properties27-29. Similarly, in these h-WO3 / m-WO3 nanoscaffolds, the coherent interphase boundaries along the vertical direction apply vertical strain to each phase, so the out-of-plane d-spacing is modulated by the lattice mismatch between the polymorphs along the c-axis, and by the phase fractions. With these coherent interphase boundaries in our nanoscaffolds, the compressive (tensile) strain were induced in h(m)-WO3 phases along the vertical direction, because the out-of-plane lattice parameter of h-WO3 (7.798 Å) is 1.6 % larger than that of m-WO3 (7.678 Å). In tungsten bronze systems, the intergrowth between hexagonal tungsten bronze (h-AxWO3) 10



and A-site vacant monoclinic tungsten bronze (m-AxWO3) can be made possible by doping a certain amount of intercalated ions (AxWO3 for x ≤ 0.10, A = intercalated ions), which stabilize hexagonal tungsten bronze structures26. Due to the competing driving force between the stability of h-WO3 by doping and the thermodynamic instability of metastable h-WO3, h-WO3 and m-WO3 can co-exist by forming intergrowth tungsten bronze structures. In our case, epitaxial stabilization on a substrate with hexagonal symmetry overwhelms the thermodynamic instability of h-WO3 at high growth rate, so for the first time, the h-WO3 / m-WO3 nanoscaffolds could be realized without introducing any intercalated ions (Fig. 3d).

Tunable phase fraction of h-WO3 / m-WO3 nanoscaffolds The h-WO3 phase was dominant in h-WO3 / m-WO3 nanoscaffolds at high W arrival rate (f = 10 Hz), but m-WO3 became dominant as the W arrival rate f decreased due to the thermodynamic instability of h-WO3 (Fig. 4a). As f decreased, the diffraction peak in 18.38° shifted to the right; eventually it reached 18.58° (equivalent to d-spacing = 7.678 Å) at f = 0.5 Hz (Fig. S1). All of the films had thickness of 60 nm, so epitaxial strain should be excluded as the origin of peak shift with growth rate. Instead, this peak exactly matched the (002) diffraction peak in m-WO3 films grown on (0112) Al2O3 substrates (Fig. 4a). Therefore, the proportions of m-WO3 and h-WO3 polymorphs can be tuned by adjusting the growth rate: as f decreases, thermodynamic stability appears to overwhelm substrate-induced epitaxial stabilization, and this trend increases the likelihood that the thermodynamically stable m-WO3 phase will form.7 In contrast, increasing f reduces the W arrival rate and thereby kinetically suppresses the thermodynamic transformations 11


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from h-WO3 to m-WO3 at growth temperature by limiting atomic transport, which is likely to favor the formation of metastable h-WO3 phase by epitaxial stabilization.8 The polymorph peaks could not be resolved in 2θ-ω, because (0001)h h-WO3 and (002)m m-WO3 have a similar Bragg’s peak positions, or because the signal from minor phases is too weak. Phase-resolved φ-scans (Fig. 4b) determined that the intensity of (2022)h h-WO3 decreased as f decreased and was absent at f ≤ 0.5 Hz, whereas the intensity of (202)m m-WO3 increased slightly as f decreased. The φ-scans suggest that (0001)h h-WO3 and (002)m m-WO3 coexist in all films, but that the dominance of the metastable h-WO3 phase increased with increase in f, and that the fraction of the six-fold m-WO3 phase increased with decrease in f. The dependence of the relative proportions of h-WO3 and m-WO3 phases on the growth rate f was also confirmed by the Raman spectroscopy (Fig. 4c, d). Excluding the Raman spectra from bare substrates (Fig. S5) isolated the phonon modes of WO3 nanoscaffolds with different phase fractions. The 253-cm-1, 690-cm-1, and 817-cm-1 phonon modes of ν(W-O-W) were dominant in film grown at f = 10 Hz (Fig. 4d); these modes are fingerprints of h-WO330. As f decreased, these phonon peaks gradually broadened, and the 274-cm-1 phonon mode of δ(O-W-O) and the 715-cm-1 and 806-cm-1 phonon modes of ν(O-W-O) started to appear (Fig. 4d). The Raman spectrum of the film grown at f = 0.5 Hz was similar to the spectrum of pure (002)m m-WO3 films on (0112) Al2O3 substrates without any h-WO3 peaks. Therefore, phase-resolved φ-scan HRXRD and Raman spectroscopy demonstrate that the proportions of h-WO3 / m-WO3 nanoscaffolds were tunable by precise control of growth kinetics.

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Hydrogen spillover of h-WO3 / m-WO3 nanoscaffolds Due to the difference in the octahedral connectivity and size of vacant sites in both polymorphs, their lattice response and kinetic stability are expected to be affected differently by chemical dopants: Intercalating hydrogens diffuse primarily through the vacant A-site in m-WO3, but mainly through open hexagonal c-axis tunnels in h-WO3.4 To quantify how intercalated hydrogens affect the lattice response of these h-WO3 / m-WO3 nanoscaffolds, they were hydrogenated and dehydrogenated using the in-situ “hydrogen spillover” method with nano-sized Pt catalysts24 at low temperature to avoid formation of oxygen vacancies; real-time synchrotron xray diffraction was performed during these processes. First, the response of the (001) Bragg’s peak in typical (001) m-WO3 (18.58°) on (0112)oriented Al2O3 substrates was examined by introducing H2-containing forming gas to the in situ chamber at 100 °C (Fig. 5a). In m-WO3 on (0112) Al2O3, the (001) Bragg’s peak split into two hydrogenated HxWO3 peaks in 30 min: one peak shifted to the left (18.43°) and one shifted to the right (18.73°); these changes are attributed to structural phase transition from monoclinic to a combination of cubic or tetragonal hydrogen tungsten bronze phases,31, 32 as a consequence of the non-uniform distribution of hydrogens in m-WO3 film. In our h-WO3 / m-WO3 nanoscaffolds, the Bragg’s peak of HxWO3 shifted to the right during hydrogenation, from 18.37° to 18.69° (out-of-plane lattice contraction ~ -1.7 %) at f = 10 Hz, and from 18.54° to 18.75° (out-of-plane lattice contraction ~ -1.1 % ) at f = 1 Hz, without any signature of peak splitting (Fig. 5b, c). Unlike m-WO3 (Fig. 5a), the c-axis lattice of h-WO3 phase contracts when mobile cations are introduced into hexagonal tunnels19, 13


33;

this phenomenon

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suggests that the hydrogen-induced lattice response in our WO3 nanoscaffolds was determined by the h-WO3 phase rather than by m-WO3. Characteristics of lattice response in m-WO3 (e.g., the lattice expansion by hydrogenation) were not observed in the h-WO3 / m-WO3 nanoscaffolds, even when it had a relatively high fraction of m-WO3 (i.e., h-WO3 / m-WO3 nanoscaffolds grown at f = 1 Hz, Fig. 5c). As shown in HAADF images (Figs. 3b-c, S4), h-WO3 and m-WO3 phases are strongly bonded by coherent interphase boundaries, and edge dislocations toward YSZ substrates exist in the m-WO3 phase; these are the reasons that the unusual lattice response of m-WO3 in the WO3 nanoscaffold strongly depends on the lattice response of h-WO3. This structural analysis of HxWO3 nanoscaffolds with coherent interphase boundaries suggests that the hydrogen-induced lattice responses of h-WO3 / m-WO3 nanoscaffolds were dominated by the hexagonal phase. During the dehydrogenation process at room temperature, hydrogen started to escape from the WO3 lattice when the dry air was introduced into the chamber. In monoclinic HxWO3 films on (0112) Al2O3 substrates, the split peaks in 18.43° and 18.73° merged to reform the original mWO3 peaks (18.58°). However, in h-WO3 / m-WO3 nanoscaffolds, the contracted lattice in HxWO3 nanoscaffolds expanded for the first 2 h, and the peaks never returned to their original positions (Fig. 5d); this result indicates that at room temperature, hydrogens bonded more strongly to the lattice in the h-WO3 phase than in the m-WO3 phase. These strong interactions between hydrogens and h-WO3 lattice occur because the intercalated h-WO3 phase is energetically more favorable than the pristine h-WO3 phase. In particular, cations that intercalate into trigonal cavity sites are bound too tightly to deintercalated from h-WO3 phase20, 21. Therefore, the energy required to extract hydrogen from a trigonal cavity of h-WO3 phase is higher than the energy to insert 14



hydrogen.

Conclusion We have shown that the h-WO3 / m-WO3 nanoscaffolds could be fabricated by delicately controlling the balance between artificially-stabilized h-WO3 phase and thermodynamically-stable m-WO3 phase. The fraction of each phase was modulated by controlling the W arrival rate at growth temperature of 500 °C, at which the m-WO3 phase is thermodynamically favorable. As a result of the small lattice mismatch along the out-of-plane direction, atomically abrupt and coherent interfaces formed spontaneously between h-WO3 and m-WO3 along the vertical direction; these interfaces led to unexpected growth of six-fold symmetrical m-WO3 phases. Due to the strong interactions between hydrogens and h-WO3 lattice from the energetically-favorable intercalated hWO3 phase, the lattice response is dominated by h-WO3, so these h-WO3 / m-WO3 nanoscaffolds showed irreversible lattice response by hydrogen intercalation. These epitaxial nanoscaffolds with different fractions of the polymorphs may allow systematical study of intercalation chemistry and its applications.

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Acknowledgement We acknowledge support for this work by the Basic Science Research Program (2017R1A2B2007819) and Creative Materials Discovery Program (2018M3D1A1058997) through the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT, and by Ministry of Trade, Industry and Energy (10076608).

Supporting Information The Supporting Information is available free of charge on the ACS Publications website.

Competing financial interest The authors declare no competing financial interest.

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Figure Captions Figure 1 | h-WO3 / m-WO3 nanoscaffold assisted by epitaxial stabilization. (a) Schematics of h-WO3 / m-WO3 nanoscaffold formation during pulsed laser deposition. The phase fraction could be modulated by W arrival rate during the growth. Crystal structure and lattice parameters of (b) m-WO3, (c) h-WO3 and (d) Yttria-stabilized Zirconia (YSZ). The metastable h-WO3 phases exhibit an empty intercrystalline tunnel along the c-axis (diameter ~ 5.36 Å). In spite of different crystal symmetry along the in-plane direction, m-WO3 (𝑑001 = 7.678 Å) and h-WO3 (2 × 𝑑0001 = 7.798 Å) show small lattice mismatch along the out-of-plane direction ([0001]h || [001]m).

Figure 2 | Epitaxial stabilization of h-WO3 / m-WO3 nanoscaffold (f = 10 Hz) on (111) YSZ substrate. (a) Symmetric 2θ-ω scan of h-WO3 / m-WO3 nanoscaffold (f = 10 Hz) on (111) YSZ substrates. Only one Bragg’s peak without other planes represents preferentially oriented growth. (b) Skew-symmetric φ-scans of h-WO3 / m-WO3 nanoscaffolds (f = 10 Hz). Based on the φ-scans, the epitaxial relation among h-WO3 (red) m-WO3 (blue), and YSZ (black) was confirmed as (0001)h [1000]h h-WO3 || (001)m [100]m m-WO3 || (111)YSZ [110]YSZ YSZ.

Figure 3 | Atomic-resolution analysis of WO3 nanoscaffolds characterized by electron microscopy. (a) Left: HRTEM image of h-WO3 / m-WO3 nanoscaffolds (f = 10 Hz) along the [11 0]YSZ zone axis. (scale bar = 20 nm) Right: selected area diffraction pattern (SADP) from the area in the left image. Blue dotted line: pseudocubic-like diffraction pattern of m-WO3; red dotted line: the tetragonal diffraction pattern of h-WO3.; cyan: diffraction pattern of YSZ substrate. (b) Magnified HAADF-STEM image along the [110]YSZ zone axis near the interphase boundaries. Scale bar: 2 nm. Abrupt vertical interfaces occur between m-WO3 phase (region expressed with purple) and h-WO3 phase (region with grayscale). In the h-WO3 phase, the strong and weak atomic column density alternates, whereas in the m-WO3 phase, similar atomic column density repeats. hWO3 phase was coherently bonded to YSZ substrate, but the m-WO3 phase shows edge dislocation 22



toward YSZ substrate due to the lattice mismatch between m-WO3 and YSZ. (c) Magnified HAADF-STEM (top panel) and ABF-STEM (bottom panel) images near the interphase boundaries between m-WO3 phase (purple region) and h-WO3 phase (grayscale region); scale bar = 1 nm. The simulated HAADF-STEM and ABF-STEM images were superimposed on the top and bottom panels, respectively. Atomic contrast is proportional to the atomic number, so HAADF-STEM detected only W atoms, whereas ABF-STEM detected both W and O. (d) Crystallographical illustrations of WO3 nanoscaffolds along the [111]YSZ projection, out-of-plane direction of the thin film. In h-WO3, strong and a weak atomic contrasts alternate along the [110]YSZ zone axis.

Figure 4 | Tunable phase fraction of h-WO3 / m-WO3 nanoscaffolds by changing growth rate. (a) Symmetric 2θ-ω scans of h-WO3 / m-WO3 nanoscaffolds with different laser frequency grown on (111) YSZ substrates. (Symmetric 2θ-ω scans of pure m-WO3 films (black) grown on (0112) Al2O3 substrate are shown for comparison) As f decreases, the diffraction peak of the h-WO3 / mWO3 nanoscaffolds shifts to right and the lattice constants of nanoscaffolds decrease (Fig. S1); these trends indicate that thermodynamic stability appears to overwhelm substrate-induced epitaxial stabilization; the result is that decrease in f yields an increase in the likelihood that thermodynamically stable m-WO3 phase will form. (b) Magnified skew-symmetric φ-scans around h-WO3 (2022)h and m-WO3 (202)m peaks of h-WO3 / m-WO3 nanoscaffolds grown with different 0.5 ≤ f ≤ 10 Hz. These scans indicate that two polymorphs coexist in all films, but that the metastable h-WO3 phase became increasingly dominant as f increased, whereas the fraction of the six-fold m-WO3 phase increased as f decreased. (c) Raman spectra of h-WO3 / m-WO3 nanoscaffolds and m-WO3 films grown on (111) YSZ and (0112) Al2O3 substrates and (d) magnified Raman spectra near the fingerprints of h-WO3 (red dotted line) and m-WO3 (black dotted line) phases. Results confirm that the relative fractions of h-WO3 and m-WO3 in the nanoscaffolds were tunable by precise control of growth kinetics.

Figure 5 | In situ structural modulation of h-WO3 / m-WO3 nanoscaffolds during 23


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hydrogenation and dehydrogenation. In situ symmetric 2θ-ω scans of (a) m-WO3 films on (01 12) Al2O3 and h-WO3 / m-WO3 nanoscaffolds grown with (b) f = 10 Hz, (c) f = 1 Hz on YSZ substrates. (d) Change in out-of-plane lattice parameters of three WO3 films during hydrogenation and dehydrogenation. t = 0 represents the starting point of dehydrogenation from hydrogenation.

24



For Table of Contents Use Only

Nanoscaffold WO3 by kinetically-controlled polymorphism

Jaeseoung Park1, HyojinYoon1, Si-Young Choi1,a) and Junwoo Son1,b)

1 Department

of Materials Science and Engineering, Pohang University of Science and Technology (POSTECH), Pohang 37673, Republic of Korea

TOC graphic

Synopsis We demonstrate an h-WO3 / m-WO3 nanoscaffold thin films by balancing the epitaxial stabilization of h-WO3 on the lattice-matched substrates against the thermodynamic instability of metastable h-WO3. The WO3 nanoscaffold is vertically aligned with coherent interphase boundaries between h-WO3 and m-WO3.

25


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Fig.1



Fig.2


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Fig.3



Fig.4


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Fig.5


Nanoscaffold WO3 by Kinetically Controlled Polymorphism - Crystal

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