Facet-Dependent in situ Growth of Nanoparticles in Epitaxial Thin Films

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Facet-Dependent in situ Growth of Nanoparticles in Epitaxial Thin Films: The Role of Interfacial Energy Kun Joong Kim, Hyeon Han, Thomas Defferriere, Daseob Yoon, Suenhyoeng Na, Sun Jae Kim, Amir Masoud Dayaghi, Junwoo Son, Tae-Sik Oh, Hyun Myung Jang, and Gyeong Man Choi J. Am. Chem. Soc., Just Accepted Manuscript • DOI: 10.1021/jacs.9b02283 • Publication Date (Web): 16 Apr 2019 Downloaded from http://pubs.acs.org on April 16, 2019

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Journal of the American Chemical Society

Facet-Dependent in situ Growth of Nanoparticles in Epitaxial Thin Films: The Role of Interfacial Energy

Kun Joong Kima‡†, Hyeon Hana§†, Thomas Defferriereb, Daseob Yoona, Suenhyoeng Naa, Sun Jae Kima, Amir Masoud Dayaghia, Junwoo Sona, Tae-Sik Ohc, Hyun Myung Janga#* and Gyeong Man Choia,d*

aDepartment

of Materials Science & Engineering / Fuel Cell Research Center, Pohang University of Science and Technology (POSTECH), Pohang 37673, Republic of Korea bDepartment

of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA, 02139, USA cDepartment d1FCell

of Chemical Engineering, Auburn University, Auburn, AL 36849, USA

Inc., Pohang 37673, Republic of Korea

‡Present

address: Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA, 02139, USA § Present

address: Max Planck Institute of Microstructure Physics, Weinberg 2, 06120 Halle (Saale),

Germany #Present

address: Research Institute of Advanced Materials, Seoul National University, Seoul 08826, Republic of Korea

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Abstract Nucleation of nanoparticles using exsolution phenomenon is a promising pathway to design durable and active materials for catalysis and renewable energy. Here, we focus on the impact of surface orientation of the host lattice on the nucleation dynamics to resolve questions with regards to “preferential nucleation sites”. For this, we carried out a systematic model study on three differently oriented perovskite thin films. Remarkably, in contrast to the previous bulk powder-based study suggesting that the (110)-surface is a preferred plane for exsolution, we identify that other planes such as (001) and (111)-facets also reveal vigorous exsolution. Moreover, particle size and surface coverage vary significantly depending on the surface orientation. Exsolution of (111)-oriented film shows the largest number of particles, the smallest particle size, the deepest embedment, and the smallest and most uniform interparticle distance among the oriented films. Based on classic nucleation theory, we elucidate that the differences in interfacial energies as a function of substrate orientation play a crucial role in controlling the distinct morphology and nucleation behavior of exsolved nanoparticles. Our finding suggests new design principles for tunable solid-state catalyst or nano-scale metal decoration. Keyword; perovskite oxide; catalysts; orientation; particle size distribution; nucleation theory

■ INTRODUCTION In perovskite materials, exsolution (i.e., in situ growth of nanoparticles) can be used to uniformly cover the oxide surface with nano-scaled catalytic metals. The exsolved nanoparticle catalysts can be cost effectively processed while possessing high coarsening and coking resistance thus making them very promising for high performance and stable solid-state catalyst especially for their use in automotive emission gas control,1,2 high temperature fuel cell,3,4 water splitting device5 and chemical looping process6. The performance of such devices could be heavily enhanced by modulating the size and number of catalysts. Several material families have attracted much interest, such as single perovskite structures based on LaFeO3,1,2,7,8 CaTiO3,1,8 LaCrO3,3,9 (La,Sr)TiO3,4,10–14 (La,Sr)ScO3,15 or Sr(Ti)FeO3,16–18 and double perovskite structures such as Sr2Fe2O6,19,20 and CoFeAlOx spinel,6 all doped with catalytic metal cations such as Rh, Pd, Pt, Ru, Ni, Co or Fe. The key requirements to promote exsolution are that theses catalyst elements be soluble in the oxide at high oxygen partial pressure (oxidizing conditions), and that the catalyst elements have a relatively low free energy of 2 ACS Paragon Plus Environment

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oxide formation to diffuse out and precipitates as nano-scale metallic particles upon lattice reduction. The fundamental condition of exsolution of B-site cation in ABO3 perovskite oxide is quite different from A-site cation segregation on surfaces which generally leads to degradation of the stability of solid oxide fuel cells (SOFCs) cathode. The segregation of particular oxide, e.g. SrO, induces a slow rate of oxygen reduction reaction and serves as a harmful barrier to implementing high-performance cathodes in oxidizing atmosphere.21-25 In contrast, the exsolution of B-site cation, occurred in highly reducing atmosphere, provides catalytically active metal nanoparticles and can be further applied to energy conversion devices. Various methodologies to trigger or accelerate the process, as well as elucidating the local surface evolution during exsolution phenomenon have been investigated. For example, the introduction of Asite vacancies destabilize the perovskite unit cell to form locally-isolated B-O bonding from the main perovskite framework, which facilitates the promotion of B-site cationic exsolution.26 Voltage-driven reduction accelerates exsolution kinetics by up to two orders of magnitude, which cannot be achieved by conventional reduction in hydrogen on the same time scale.4 Numerical simulations based on strain-field modeling to analyze surface evolution during exsolution process suggested that when reduced, Ni particles deform the surrounding oxide and induce an elastic strain in the system, where the interplay between surface free energy and the strain energy associated with the metal particles leads to particle-in-a-pit morphology.27 Aforementioned studies have reinforced the understanding of the exsolution process. However, it is also important to investigate the impact of key material parameters on controlling the size and surface coverage of the catalysts, which could ultimately govern the performance of the devices. In classical nucleation theory the driving force for nucleation is the cation supersaturation of the host lattice during the reduction. While supersaturation may drive the nucleation process and also determine some aspects of the morphological outcome, it is not the only factor to consider regarding to the particle size distribution achieved. The particle size distribution, which is an outcome of the nucleation process, may be furthermore influenced by surface and interfacial energy contributions of the nucleated particles as well as local strain fields in the host matrix. Several reports already confirmed that particle size and coverage of exsolved catalyst on the perovskite surface can be modulated by controlling the degree of supersaturation via extrinsic operating parameters such as the reduction temperature and time3,9,20,28 or oxygen partial pressure of reducing gas15 or both29. For titanate polycrystalline bulk systems, preferential growth of nickel nanoparticles was reported on (110) surfaces over (001) and (111) surfaces.20,30 This interesting fact suggests that surface orientation 3 ACS Paragon Plus Environment


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could play an important role in modulating the nucleation energy barrier, which in turn, controls the particle size distribution and surface coverage. However, polycrystalline sintered bodies30 have typically been used for exsolution studies making it difficult to deconvolute the contribution of surface, strain and interfacial energies. Indeed, simple model experiments to study the impact of surface orientation on the exsolution behavior is missing and hence the question of how particle size and coverage can be manipulated remains unexplored. To resolve the question, in this work, we report exsolution of Nickel nanoparticles on (100)-, (110)- and (111)-oriented strontium titanate perovskite (La0.2Sr0.7Ti0.9Ni0.1O3-δ) thin films. Distinct morphological outcome and chemical features of the nanoparticle shows that exsolution on the (111) facet reveals the largest number of particles (195 m-2), the smallest particle size (20 ± 7 nm), among the (100)-, (110)- and (111)-oriented films. Major energy contributions to the exsolution process controlling the particle size distribution and particle contact angle are discussed based on the classical nucleation theory. ■ RESULTS We fabricated epitaxial La0.2Sr0.7Ti0.9Ni0.1O3-δ (LSTN) thin films by pulsed laser deposition (PLD) on the (001), (110), and (111) SrTiO3 (STO) substrates with a thickness of ~300 nm. After the deposition, the films were chemically reduced in a controlled-atmosphere furnace (dry H2, 900 °C for 12 h) to exsolve the Ni particles to the surface. The surface of as-reduced LSTN films were investigated by atomic force microscopy (AFM) to observe the effect of reduction on microstructural evolution. First, we observe that metallic secondary phases have exsolved from the film to its surface as shown in Figure 1 (elemental confirmation is provided in the next section). AFM reveals that a root mean square surface roughness (rRMS) decreases from 4.8 nm, 3.1 nm to 1.4 nm for (001)- (Figure. 1a, b, c), (110)- (Figure 1d, e, f) and (111)-oriented films (Figure 1g, h, i), respectively. The rRMS for as-deposited films also decreases from 0.378 nm, 0.310 nm to 0.285 nm for the (001)-, (110)- and (111)-oriented LSTN films, respectively (Figure S1). We now probe the role of LSTN orientation on the exsolution of Ni in terms of the size and distribution (Table 1 and Figure S2). The number of particles and particle width/height were obtained from the AFM results in Figure 1, and the interparticle distance was calculated based on the SEM images. The average particle width (w) decreases from 40 ± 17 nm, 31 ± 9 nm to 20 ± 7 nm for the (001)-, (110)-, and (111)-oriented LSTN films, respectively, and the average particle height (h) decreases from 13 ± 6.0 nm, 7.4 ± 3.7 nm to 3.7 ± 1.4 nm for the (001)-, (110)-, and (111)-oriented LSTN films, respectively. Exsolved Ni particles show a spherical-cap shape with average aspect ratio h/w of ~0.33, ~0.24 and ~0.19 for (001)-, (110)-, 4 ACS Paragon Plus Environment

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and (111)-oriented films, respectively. The number of particles per μm2 area also vary: 52 μm-2 for (001)-oriented film, 111 μm-2 for (110)-oriented film, and 195 μm-2 for (111)-oriented film. As well as the particle size, average interparticle distance (IPD) between two adjacent particles can be also an important factor in catalytic applications.31,32 The IPD of each film decreases in the order of (001), (110), and (111) samples: 77 ± 50 nm, 49 ± 40 nm, and 36 ± 18 nm, respectively. The small IPD along with the smallest particle size of the (111) sample could be most advantageous in terms of catalytic activity, since it is not too small spacings to cause negative effects such as shielding effect or diffusion effect.31 Summarizing the results of AFM and SEM, exsolution of (111)-oriented film shows the smallest particles size, the largest number of particles, and the smallest and most uniform IPD among the three oriented films. Another notable thing is that all the oriented thin films reveal vigorous exsolution. These results are in contrast with the bulk ceramic systems which have shown preferential exsolution behavior along (110) facet.30 To obtain further insight on the exsolution phenomenon in the film, scanning transmission electron microscopes (STEM) was used for a representative (001)-oriented film after reduction (Figure 2). Selected area electron diffraction (SAED) patterns for the LSTN film (Figure 2a) and for the STO substrate (Figure 2b) along the zone axis [100] show no ring patterns, and the calculated inter-planar d-spacings for film and substrate are almost the same (Table S1), which confirm unequivocally epitaxial orientation relationship of LSTN (001)/STO (001). Cross-sectional image shows that the 300 nm-thick film reveals a dense microstructure with a number of local region of bright contrast (Figure 2c, d). Elemental profile confirmed that these are mainly composed of metallic Ni (Figure 2e, f) other than A-site cation segregation such as SrO (Figure S3). Crystallographic information of the films before and after exsolution is important to precisely define surface orientation as well as to understand the possible contribution of lattice strain on the observed tunable catalyst in this system. θ−2θ XRD patterns show that both “as-deposited” and “as-reduced” films are epitaxial, thus highly oriented on the three STO substrates (Figure 3a, b, c); full-range XRD scans and in-plane (IP) alignments of the film stacks were investigated to confirm the orientation (Figure S4). The  scans of both as-deposited and as-reduced films reveal identical symmetry with the substrates, which means that the epitaxial growth of the films on the substrates is cube-on-cube and agree with aforementioned SAED patterns (Figure 2a, b). The diffraction peaks of all asdeposited LSTN film are located in lower 2θ values compared with LSTN bulk and STO substrate, implying that the out-of-plane (OOP) lattice spacing of the LSTN films are larger than that of LSTN bulk and STO substrate. Remarkably, OOP lattice spacing values of as-reduced films are decreased 5 ACS Paragon Plus Environment


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in comparison with those of the as-deposited films: 3.938 Å → 3.907 Å, 2.771 Å → 2.767 Å, and 2.260 Å → 2.257 Å for (001), (110), and (111)-oriented films, respectively. To further identify the change of both OOP and IP lattice parameters, we conducted reciprocal space mapping (RSM). In as-deposited films, the OOP lattice expanded, so the reciprocal lattice points of each film are at a smaller value on the Qz axis than the point of the substrate (Figure 3d, e, f); these results correspond to the θ−2θ XRD results. In contrast, the reciprocal lattice point of as-deposited LSTN films is located nearly the same with that of STO along the Qx axis in each map. A sharp film peak was aligned with the STO substrate peak; this correspondence indicates that the film has strong anisotropic strain without noticeable relaxation.33 After reduction, unit cells shrink along the OOP direction, but IP lattice parameters show a negligible change: 3.904 Å → 3.906 Å, 2.762 Å → 2.762 Å and 2.758 Å → 2.760 Å for (001), (110) and (111)-oriented films, respectively (Figure 3g, h, i). That is, the films are rigidly clamped on the substrates even after the strong reduction. The as-deposited films have larger unit cell volumes than the bulk polycrystalline powder, as reported in other oxide film deposited by PLD i.e. (La,Sr)CoO3 or ceria-based films.34,35 The structure of as-deposited (001)-oriented film is tetragonal, as-deposited (110)-oriented film is double-cell orthorhombic or equivalently limiting monoclinic, and as-deposited (111)-oriented film is rhombohedral (Figure S5).36 The XRD analysis reveals two notable points: (i) after reduction, lattice spacing of the films approach bulk values similar to macro-crystalline pellets (Figure 3a, b, c, Figure S6) while the films retain their crystallinity (Figure 3 and S4). The result means that initially stretched as-deposited films come back to the bulk unit cell dimension upon reduction. (ii) the change of OOP lattice spacing is larger in the (001)-oriented film (−0.77 %) when compared to the (110)-oriented (-0.16 %) and (111)-oriented films (−0.14 %). To probe surface chemistry of LSTN films, X-ray photoelectron spectroscopy (XPS) patterns were obtained from (001)-, and (111)-oriented films (Figure 4). A summary of numerical data for La, Sr, Ti and Ni (effective peak area and respective atomic ratio) before and after reduction, respectively, is provided (Table S2). In reducing atmospheres, A and B site cations in AB1-xMxO3 perovskite are typically maintained in oxidized states while M such as transition metal is easily reduced and exsolved (not necessarily all to the metallic state) due to its limited thermodynamic stability.15 As a result, the electro-valence values of Ni (3p) and Ti (3s and 2p) XPS peaks changed after reduction. For simplicity of calculation, we considered only Ni 3p 3/2 and Ti 2p 3/2 peaks. In (001)-oriented film, 42 % of surface Ni0 was exsolved out of total Ni (Ni0 + Ni2+) (Figure 4a, b). Ti4+ was partially reduced to Ti3+, leading to Ti3+/(Ti3++Ti4+) ~1.05 % (Ti 3s peak) and ~1.1 % (Ti 2p peak) (Figure 4c, d); the amount of reduced Ti3+ was much lower in the film than the bulk pellets (6 %).30 For the (111)-oriented film, 6 ACS Paragon Plus Environment

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24 % of surface Ni0 was exsolved out of total Ni (Ni0 + Ni2+) and Ti3+/(Ti3++Ti4+) ~0.8 % (Figure S7). The XPS data confirm that surface composition varies for every orientation, and, importantly, exsolved species are metallic Ni. ■ DISCUSSION The exsolution process is thermodynamically driven by the reduction of the total free energy of the system and subsequently accompanied by particle migration towards the surface. In the case of the bulk polycrystalline systems, there can be differences in the migration kinetics or nucleation thermodynamics due to differences in surface/interfacial energy for each crystalline plane. Thus, the outcome of the exsolution process in polycrystalline samples is typically the result of an intricate competition between differences in thermodynamic nucleation energies and differences in migration kinetics of the particle among the various crystal orientation, eventually leading to a scenario where the most favorable orientation dominates the overall exsolution behavior. That is, (110) planes have shown to be preferential exsolution surfaces in the A-site deficient bulk ceramics possibly due to the faster cation migration kinetics.30 In contrast, in the case of single crystalline thin film systems, it can be considered that there exists only a single surface through which particles can go out of the oxide matrix. Thus, the (001) and (111)-oriented films can also reveal substantial exsolution of particles toward each surface. In other words, epitaxial thin films enable tunable and distinct exsolution behaviors by just changing the substrate orientation, resulting in the distinct particle morphology with the differently oriented films. To identify the distinct exsolution behavior near the surface in each oriented thin film, we used the homogeneous nucleation model as follows: ∆𝐺ℎ𝑜𝑚 = ∆𝐺𝑣𝑜𝑙 + ∆𝐺𝜀1 + ∆𝐺𝜀2 + ∆𝐺𝑑𝑖𝑓𝑓 + ∆𝐺𝑖𝑛𝑡 + ∆𝐺𝑠𝑢𝑟 (1) which can be rewritten as ∆𝐺ℎ𝑜𝑚 = 𝑉(∆𝑔𝑣𝑜𝑙 + ∆𝑔𝜀1 + ∆𝑔𝜀2 + ∆𝑔𝑑𝑖𝑓𝑓) + 𝐴𝑖𝑛𝑡𝛾𝑖𝑛𝑡 + 𝐴𝑠𝑢𝑟 𝛾𝑠𝑢𝑟 (1𝑎) 4 ∆𝐺ℎ𝑜𝑚 = 𝜋𝑟3(∆𝑔𝑣𝑜𝑙 + ∆𝑔𝜀1 + ∆𝑔𝜀2 + ∆𝑔𝑑𝑖𝑓𝑓) + 4(𝜋 ― 𝜃)𝑟2𝛾𝑖𝑛𝑡 + 4𝜃𝑟2 𝛾𝑠𝑢𝑟 (1𝑏) 3 where ∆𝐺 represents total energy and ∆𝑔 represents the volumetric and area normalized energy terms. ∆𝐺ℎ𝑜𝑚 is total free energy for Ni exsolution, ∆𝐺𝑣𝑜𝑙 is the volume free energy gain by nucleating 7 ACS Paragon Plus Environment


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the new phase, ∆𝐺𝜀1 is related to the strain energy between Ni and LSTN, ∆𝐺𝜀2 is effective lattice strain relaxation energy, ∆𝐺𝑑𝑖𝑓𝑓 is from the Ni nucleus diffusion towards surface, ∆𝐺𝑖𝑛𝑡 is the free energy change by introducing an interfacial area between the particle and the host matrix, ∆𝐺𝑠𝑢𝑟 is the free energy change by forming the particle on the surface. 𝛾𝑖𝑛𝑡 is the interfacial energy between the particle and the oriented perovskite lattice at the surface, 𝛾𝑠𝑢𝑟 is the surface energy of the particle, 𝜃 is the particle contact angle on the surface, and 𝑟 is the particle radius. The ∆𝐺𝑣𝑜𝑙 and ∆𝐺𝜀2 are negative terms, while the others are positive terms. A constant steady-state nucleation rate (𝑁) and critical particle radius (𝑟 ∗ ) for a given T and driving force are derived by differentiating ∆𝐺ℎ𝑜𝑚 with respect to 𝑟 and equaling zero:37

∗

𝑟 =―

∗ ∆𝐺ℎ𝑜𝑚

2[(𝜋 ― 𝜃)𝛾𝑖𝑛𝑡 + 𝜃𝛾𝑠𝑢𝑟] 𝜋(∆𝑔𝑣𝑜𝑙 + ∆𝑔𝜀1 + ∆𝑔𝜀2 + ∆𝑔𝑑𝑖𝑓𝑓)

(2)

[(𝜋 ― 𝜃)𝛾𝑖𝑛𝑡 + 𝜃𝛾𝑠𝑢𝑟]3 4 16 2 ∗ (𝑟 ) = [(𝜋 ― 𝜃)𝛾𝑖𝑛𝑡 + 𝜃𝛾𝑠𝑢𝑟](𝑟 ) = (3) 3 3 [𝜋(∆𝑔𝑣𝑜𝑙 + ∆𝑔𝜀1 + ∆𝑔𝜀2 + ∆𝑔𝑑𝑖𝑓𝑓)]2 ∗

𝑁 ∝ 𝑒𝑥𝑝

[

∗ ―∆𝐺ℎ𝑜𝑚 (𝑟 ∗ )

𝑘𝑇

]

(4)

Equation (2) is the critical nucleation size which represents the nuclei size above which particles are stable and can grow. Equation (3) is the critical energy barrier to nucleation to achieve a particle of critical nuclei size 𝑟 ∗ . Equation (4) is the nucleation rate of those particles. The values of the terms in Equation (1) are system-dependent. Based on our experimental data, we believe that the main parameters affecting the exsolution process that need to be considered when changing the crystal orientation from (001) to (111) are the lattice strain relaxation energy (∆ 𝐺𝜀2) along OOP and interfacial energy (∆𝐺𝑖𝑛𝑡) between the Ni particles and the LSTN host. Since the experiments have been conducted at isothermal conditions and fixed oxygen partial pressure (implying similar amounts of supersaturation) the two free energy terms from bulk reduction (∆𝐺𝑣𝑜𝑙) and diffusivity (∆𝐺𝑑𝑖𝑓𝑓) might be the same with respect to orientation. When we look carefully at the morphological profile (Figure 5) of the particles based on AFM measurements (Figure 1) with different sample orientation, as for example of (001)- and (111)-oriented sample, they show spherical Ni 8 ACS Paragon Plus Environment

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particles with no sign of preferred facet formation. However, Ni particles found on the surface had different average depths into the oxide film depending on the film orientation. The particle depth refers to the Young-Dupre (Winterbottom) framework38 where the particles minimize their total energy by maximizing the surface with the lowest energy (i.e., solid state equivalent to wetting of liquid droplets on the substrate). In contrary to liquid nucleation where the particle can spread in order to wet, in solid state, the particle varies the extent of embedment into the matrix to wet. Ni particles exsolved on (111)-oriented film are more deeply embedded in the LSTN host compared to the (001)-orientation, which may imply that oxide film orientation induces metal-oxide interface energy difference. Based on these experimental observations, we can conclude that both interfacial energy (∆𝐺𝑖𝑛𝑡) and lattice strain relaxation energy (∆𝐺𝜀2) vary for each oriented film and thus might play a critical role in determining the particle size distribution of the exsolved particles. Therefore, equation (2) and (3) can be simplified and rewritten by considering the dominant terms only:

𝑟∗ ∝ ―

∗ ∆𝐺ℎ𝑜𝑚

(𝜋 ― 𝜃)𝛾𝑖𝑛𝑡 ∆𝑔𝑣𝑜𝑙 + ∆𝑔𝜀2

(𝑟 ) ∝ ∗

(5)

[(𝜋 ― 𝜃)𝛾𝑖𝑛𝑡]3 (∆𝑔𝑣𝑜𝑙 + ∆𝑔𝜀2)2

(6)

We first consider the contribution of strain energy to the total free energy. Based on the XRD results in Figure 3, OOP tensile strain 𝑓 = (𝑑𝑓𝑖𝑙𝑚 ― 𝑑𝑏𝑢𝑙𝑘) 𝑑𝑏𝑢𝑙𝑘 can be calculated: 0.76, 0.27, and 0.17 % in asdeposited (001)-, (110)- and (111)-oriented films, respectively (Table S3). The different amounts of the strain in differently oriented films can be explained by looking at the Poisson`s ratio. Using the elastic compliances values for S11, S12, and S44 for undoped-SrTiO339 and direction cosines for the plane direction, elastic constant E of 265, 287 and 295 GPa were calculated for (001)-, (110)- and (111)orientations (see Supplementary Note); i.e., the largest OOP strain occurs in (001)-oriented film and the smallest OOP strain occurs in (111)-oriented film. If the strain dominates the total free energy, then according to Equations (4-6), for the most severely pre-strained (001)-oriented film, we would obtain a smaller particle size and higher nucleation rate.40 However, this is inconsistent with our experimental observation (Figure 1). We find smaller particle sizes and higher surface area coverages for (111)oriented films. Unlike the effect of strain energy, the interfacial energy reveals a great influence on the total free energy. Based on AFM results in Table 1, the embedded particle radii can be expressed as 9 ACS Paragon Plus Environment


ℎ2 +

𝑟=

(𝑤 2) 2ℎ

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2

, and their effective contact angle can be defined as 𝜃 = 𝑐𝑜𝑠 ―1

( ). The calculated 𝑟―ℎ 𝑟

embedded particle radii (𝑟) are 22, 20, and 15 nm, and contact angles (𝜃) reveal 66, 51, and 41° for (001)-, (110)-, and (111)-oriented films, respectively. The obvious differences in the particle contact angles and the corresponding interfacial area (Figure 5) in each oriented film provide strong evidence of distinct interfacial energies (𝛾𝑖𝑛𝑡) in each film. That is, the interfacial energy (𝛾𝑖𝑛𝑡) in (001)-film is relatively higher than in (111)-film because the thin film projection area-normalized Ni-LSTN interfacial area (𝐴𝑖𝑛𝑡) in contact with the LSTN host is smaller in the (001)-film than in the (111)-film (Figure 5). For higher 𝛾𝑖𝑛𝑡, the total nucleation energy barrier (equation 6) as well as the critical nucleation size (equation 5) increases. Moreover, the nucleation rate (𝑁), i.e., the particle population density (equation 4), decreases as a result of increased nucleation barrier which seems to be consistent with our observation. Based on the model described above, therefore, we judge that the nucleation behavior and the particle morphology in this system is mainly governed by the variation in the interfacial energy as a function of surface orientation. The present results suggest that surface orientation and the interfacial energy of the LSTN films to the exsolved particles strongly influence the dynamics of the Ni exsolution, which are important parameters for tunable catalyst preparation. We believe our model study provides useful insight to modulate catalyst design on the film surfaces and it would be interesting by combining other various methods that have been developed to accelerate exsolution process such as control of stoichiometry (defect engineering)26, electrochemical poling4 or control of thermodynamic conditions (reduction Po2 and temperature).15,41,42 ■ CONCLUSIONS We have studied the impact of the surface orientation of the host matrix on exsolution of nanoparticles in three differently oriented thin films, resulting in distinct particle morphologies and surface coverages. A classic homogenous nucleation model was considered to demonstrate the main parameters controlling the thermodynamics of exsolution on the basis of changes in the lattice strain and the exsolved particle morphology. The proposed model describes the critical role of the interfacial energy on the exsolution process. Ni particles exsolved from the (001)-oriented surface have shown the highest interfacial energy, resulting in the largest particle size with the least surface coverage among the (001), (110), (111)-oriented perovskites and with the smallest degree of embedment into the substrate. On the contrary, due to the lowest interfacial energy in the (111)-oriented surface, it 10 ACS Paragon Plus Environment

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has shown to achieve the smallest particles size, the highest surface coverage, the deepest embedment, and the smallest and most uniform interparticle distance, which could be a good candidate for catalytic applications. This study suggests potential to tune exsolution of metal particles on the hetero-epitaxial oxide films by just altering substrate orientations and provides a route to highly active catalysts that are sufficiently stable as an application in energy-conversion and storage systems.

■ EXPERIMENTAL DETAILS Thin Film Fabrication: Pulsed laser deposition (PLD) method was used to deposit (001), (110) and (111)-oriented La0.2Sr0.7Ti0.9Ni0.1O3- (LSTN) films on single-crystalline SrTiO3 (001), (110) and (111) substrates (MTI Corporation), respectively. Facet-dependent epitaxial growth can be possible due to the small lattice mismatch (-0.5 %) between LSTN film (a = 3.927 Å) and SrTiO3 substrate (a = 3.906 Å).40 During deposition, laser energy density was 1.5 J cm-2 and the repetition rate was 5 Hz; the substrate was maintained at 800 °C with the Po2 at ~50 mTorr. For PLD target, LSTN powder was synthesized using a solid-state reaction method according to Ref.43, calcined at 1300 °C, die-pressed to form a pellet and sintered at 1500 °C for 5 h of isothermal hold in air. All film samples analyzed in the present work were from a single target. Sample Reduction: Dry H2 supplied directly from the gas cylinder served as the reducing agent. Reducing gas flow started at RT and stopped after the film deposition and the sample had cooled to RT. The samples were reduced at 900 °C for 12 h at a heating rate ~3 °C min-1 and cooling rate ~5 °C min-1. Characterization: We performed XRD structural analysis to confirm the presence of the perovskite phase. The phase of LSTN bulk powder was determined using XRD (Model D/MAX-2500/PC, RIGAKU, Japan) employing Cu-Kα radiation ( Å ), operated at 200 mA and 40 kV. The lattice parameter of bulk powder was calculated using a program (“Unit Cell”, Department of Earth Sciences, University of Cambridge, UK). Silicon was used as a reference for calibration of the peak position. Thin films were characterized using a high-resolution XRD (D8 Discover, Bruker) under Cu Kα radiation; and reciprocal space mapping was conducted using the same device. The microstructure of the samples was examined using a field-emission scanning electron microscope (XL30S FEG, Philips Electron Optics B.V., Netherlands). Surface profiles of the exsolved particles were collected by AFM (Park systems, XE-100, South Korea) in non-contact mode. The 11 ACS Paragon Plus Environment


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microstructure of the LSTN films on STO substrate was examined using scanning transmission electron microscopy (JEOLJEM- 2100F, JEOL with a Cs-corrector). The interparticle distance was calculated for an area of 1 μm2 in the SEM images using ImageJ software. Elemental composition and valence near the surface were measured using XPS (AXIS Ultra DLD, Kratos. Inc). The spectra were calibrated based on the C 1s peak from adventitious carbon. Quantification was performed based on the area of peaks of interest (La 4d, Sr 3d, Ti 2p and 3s, Ni 3p) after deconvolution of each peak using a linear-type background subtraction using XPSPEAK software.

■ ASSOCIATED CONTENT * Supporting Information The Supporting Information is available: AFM, IPD and SAED patterns analyses, XRD Phi scans, strained structure of the films, θ-2θ XRD scan of the bulk, XPS, and OOP strain and elastic constant calculation ■ AUTHOR INFORMATION Corresponding Authors *[email protected] *[email protected] Author Contributions †K.

J. Kim and H. Han contributed equally to this work.

Notes The authors declare no competing financial interest. ■ ACKNOWLEDGMENTS The authors acknowledge Dr. J. W. Won for his help in calculating the elastic constant of SrTiO3. Mr. S. B. Yoo, J. S. Park and C. D. Oh are thanked for discussion and experiments for AFM and Raman spectroscopy. We thank also Dr. G. H. Lee, J. C. Rosillo and J. M. L. Rupp for the comments on the manuscript. This work was supported by the National Research Foundation (NRF) Grant funded by the Korean Government (MSIP) (Grant No. 2016R 1D1A1B 03933253).


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Figure 1. AFM and SEM top-view images of as-reduced LSTN films for (a, b) (001)-oriented film, (d, e) (110)-oriented film, (g, h) (111)-oriented film. Scale bars in Figure 1b, e, h: 1 m, 200 nm for insets. Figure 1c, f, i show schematics of the surface of LSTN films decorated with exsolved Ni nanoparticles. AFM image area = 1 µm2. Thickness of LSTN films is ~300 nm.



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Table 1. Quantitative characteristics of exsolved Ni particles. Number of particles and particle width/height were obtained from the AFM results in Figure 1. Average and standard deviation of the inter-particle distance were calculated from the SEM images.

Film orientation

Number of particles per area [/μm2]

Particle width (w)

Particle height (h)

Interparticle distance

[nm]

[nm]

[nm]

(001)

52

40 17

13 6.0

77 50

(110)

111

31 9

7.4 3.7

49 40

(111)

195

20 7

3.7 1.4

36 18


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Figure 2. Scanning transmission electron microscopy (STEM) analysis for LSTN (001) film after reduction and thus Ni exsolution. (a, b) Selected area electron diffraction (SAED) patterns of spots in (c) marked with ‘x’; DP1 in LSTN film and DP2 in STO substrate along the zone axis [100]. (c, d) highangle annular dark field (HAADF) image. (e) Line profiles extracted from the arrow in the Figure 2d. (f) STEM-EDS map of Figure 2d. The element profile images show precipitated Ni clusters ~20 nm. Scale bar: 50 nm in 2c, 10 nm in 2d, f.



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Figure 3. Structural properties of (001)-, (110)- and (111)-oriented LSTN thin films. θ−2θ scan of the LSTN film deposited on (a) STO (001), (b) STO (110), and (c) STO (111) substrates. Semilogarithmic contour plots of the X-ray reciprocal space mappings (RSM) (d, e, f) before and (g, h, i) after reduction. The RSMs were taken for (d, g) the (001)-oriented film around the STO (103) Bragg reflections, (e, h) the (110)-oriented film around the STO (130) Bragg reflections, and (f, i) the (111)oriented film around the STO (312) Bragg reflections. After reduction, the lattice spacings of the films approach those of bulk powder. The change of OOP lattice spacing in (001)-oriented film is exceptionally larger than those of (110) and (111) films; −0.77 %, −0.16 % and −0.14 % (d-spacing contraction) for (001), (110) and (111) films, respectively. STO denotes SrTiO3; LSTN denotes La0.2Sr0.7Ti0.9Ni0.1O3-


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Figure 4. X-ray photoelectron spectroscopy (XPS) of LSTN (001) surface. Ni 3p, Ti 3s for the sample (a) before and (b) after reduction. Ti 2p for the samples (c) before and (d) after exsolution.



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Figure 5. AFM height and lateral diameter profiles (black symbol with line) of representative exsolved particle (a) (111)-oriented film, (b) (001)-oriented film. Drawing indicates exsolved and embedded Ni particles in LSTN films. Lower area-normalized interfacial area in (001)-oriented film implies higher interfacial energy.


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TOC Figure


Facet-Dependent in situ Growth of Nanoparticles in Epitaxial Thin Films

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