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Surface-orientation-dependent growth of SrRuO3 epitaxial thin films
Applied Surface Science 499 (2020) 143924
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Surface-orientation-dependent growth of SrRuO3 epitaxial thin films a
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c,d
a
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Sungmin Woo , Hyuk Choi , Seunghun Kang , Jegon Lee , Adrian David , Wilfrid Prellier , ⁎ Yunseok Kimc,d, Hyun You Kimb, Woo Seok Choia, a
Department of Physics, Sungkyunkwan University, Suwon 16419, Republic of Korea Department of Material Science and Engineering, Chungnam National University, Daejeon 34134, Republic of Korea School of Advanced Materials Science and Engineering, Sungkyunkwan University, Suwon 16419, Republic of Korea d Research Center for Advanced Materials Technology, Sungkyunkwan University, Suwon 16419, Republic of Korea e Laboratorie CRISMAT, CNRS UMR 6508, ENSICAEN, Normandie Universite, 6 Bd Marechal Juin, F-14050 Caen Cedex 4, France b c
A R T I C LE I N FO
A B S T R A C T
Keywords: Pulsed laser epitaxy Crystallographic orientation Surface energy Growth kinetics SrRuO3
The growth of epitaxial transition metal oxide thin films depends on various parameters including the substrate temperature, oxygen partial pressure, kinetics of incoming adatoms, Gibbs free energy, and surface energy. Naturally, the change in the crystallographic surface orientation with a distinctive surface energy also influences the growth rate and growth mode of the epitaxial thin films substantially. Using perovskite SrRuO3 as a model system, we studied the growth characteristics by changing the surface orientation of the SrTiO3 substrate. Employing X-ray diffraction and surface atomic and Kelvin probe force microscopy (KPFM), we observed a systematic decrease in the growth rate and a modification in the growth mode from a two-dimensional growth to a three-dimensional island growth with the change in the surface orientation from (100) to (110) to (111). A spin-polarized density functional theory calculation demonstrated the corresponding difference in the surface energy, which was also confirmed experimentally by the KPFM measurements. The difference in the surface energy could explain the observed change in the growth kinetics, based on the modified classical nucleation theory. A combinatorial method using a polycrystalline epitaxial thin film was employed to generalize our understanding of the crystallographic surface-orientation-dependent thin film growth.
1. Introduction Perovskite SrRuO3 (SRO) is a strongly correlated itinerant ferromagnet, widely employed as a metallic electrode for oxide electronics. Bulk SRO has an orthorhombic crystal structure with ferromagnetism arising below the Curie temperature (TC) of ~160 K [1]. When epitaxially grown on different perovskite substrates, it exhibits structural transition from an orthorhombic to a tetragonal structure as a function of the epitaxial strain and thickness, which influences its electronic and magnetic properties [2–6]. The crystallographic surface orientation of the epitaxial SRO thin film is also crucial for its physical behaviors. A distinctive structural distortion can be anticipated, i.e., tetragonal distortion for (100)-, monoclinic distortion for (110)-, and trigonal distortion for (111)-oriented epitaxial thin films. The difference in the structural distortion affects the electronic structure and magnetic exchange coupling, eventually modifying the itinerant nature and ferromagnetic behavior [7–10]. Therefore, precise understanding and control of the growth of the epitaxial thin films with various orientations are highly necessary.
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Indeed, the atomic-scale control of the perovskite oxide thin film growth is of prime importance for the state-of-the-art heterostructure fabrication techniques. The growth rate of the perovskite oxide thin film is determined by the combination of several parameters including the surface energies of the substrate and thin film, growth temperature and pressure, incoming atomic concentration and energy, Gibbs free energy, and nucleation and desorption energies of the atoms. The orientation dependence of the growth rate has been studied, for growth methods such as chemical vapor deposition and sputtering using substrates with different surface orientations [11–14]. However, most of the films used to analyze the growth mechanism were on the micrometer scale, such that the atomic-scale information was not accessible. In addition, the growth mode variation with the crystallographic orientation has not yet been explained in detail for a systematically controlled growth system. For better understanding of the orientationdependent growth mechanism leading to fine synthetic crystals with atomic-scale precision, studies exploiting the well-established pulsed laser epitaxy (PLE) are required. In this study, we investigated the relationship between the surface
Corresponding author. E-mail address: [email protected] (W.S. Choi).
https://doi.org/10.1016/j.apsusc.2019.143924 Received 14 June 2019; Received in revised form 28 August 2019; Accepted 9 September 2019 Available online 10 September 2019 0169-4332/ © 2019 Elsevier B.V. All rights reserved.
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energy and growth mode (and rate) of SRO epitaxial thin films by employing various surface orientations, with a specific focus on the atomic-scale precision epitaxy. We fabricated epitaxial SRO thin films on SrTiO3 (STO) substrates with (100), (110), and (111) surface orientations using PLE. Various film thicknesses in the range of a few to a hundred nanometers were investigated. Using atomic force microscopy (AFM) and Kelvin probe force microscopy (KPFM), we confirmed the growth mode variation with the surface potential of each orientation. The surface energies of the (100)-, (110)-, and (111)-oriented SRO films were calculated by a spin-polarized density functional theory (DFT) calculation, yielding results consistent with the experimentally obtained values. Based on the modified classical nucleation theory, the different unit-cell surface energies provided comprehensive understanding of the orientation-dependent growth mechanism. Furthermore, we employed epitaxial polycrystalline SRO thin films for combinatorial confirmation of the crystallographic surface-orientationdependent growth.
(3 × 3) slabs with 12 atomic layers. The topmost and bottommost layers of the model slabs were released in the optimization to calculate the energy associated with the surface relaxation. As a stoichiometric SRO slab cannot be sliced into two SRO slabs with two identical surface morphologies, we applied an average estimation method, which was also applied in our previous study [15]. Two different surface terminations were considered for each surface: RuO2 or SrO termination for (100), SrRuO or O2 termination for (110), and Ru or SrO3 termination for (111). As both surface terminations simultaneously form upon cleavage of a stoichiometric SRO slab, we initially calculated the cleavage energy of SRO, Ecleavage, and equally distributed it to both terminations. For example, the unrelaxed internal cleavage energy of the (100) surface was calculated as [21–23]. (100) Ecleavage =
1 (100) − RuO2 (100) − SrO [E + Eslab − x ΔEbulk ] 4 slab (100)−RuO2
(1)
(100)−SrO
and Eslab are the energies of the unrelaxed where Eslab RuO2- and SrO-terminated surfaces, respectively, ΔEbulk represents the formation energy of a single bulk SRO unit cell from separated SRO clusters obtained from the calculation (ΔEbulk = −5.76 eV), and x is the number of total unit cells in both slabs. For the (100), (110), and (111) surfaces, the x values were 81, 81, and 99, respectively. The factor of 1/ 4 is included as four surfaces were created upon cleavage. Based on Ecleavage and ΔEbulk, for example, the surfaces energy of the SrO-terminated (100) surface was calculated as
2. Methods 2.1. Epitaxial thin film growth and structural characterization High-quality epitaxial single-crystal and polycrystalline SRO thin films were synthesized on STO single-crystal ((100), (110), and (111)) and polycrystalline substrates using PLE at 700 °C, respectively. The substrates were treated to obtain atomically flat surfaces [15]. An excimer laser (248 nm; IPEX 864, Lightmachinery) with a fluence of 1.5 J/ cm2 and repetition rate of 5 Hz was used. The thin films were grown at an oxygen partial pressure of 100 mTorr, which results in a stoichiometric SRO thin film [3]. Each set of thin films (with the same number of laser shots) was fabricated at the same time. The atomic structures, crystal orientations, epitaxy relationships, and thicknesses of the SRO thin films were characterized using high-resolution X-ray diffraction (XRD, Malvern PANalytical, X'Pert Pro MRD XL) and electron backscattering diffraction (EBSD, JSM7000F). For the EBSD, the sample was mounted at an angle of 70° from the surface normal within a scanning electron microscope operated at 20 kV. The probe current of the aperture was 1 × 10−8 A. The Kikuchi diagrams were recorded with a beam step size of 0.5 μm. A commercial program (OIM Analysis 5.31) was used for image processing.
(100) − SrO (100) Esurface = Ecleavage +
1 (100) − SrO ∆Erelaxation 2
(2)
where ΔErelaxation(100)−SrO is the energy released upon the relaxation of the slab. 3. Results and discussion 3.1. Crystallographic surface-orientation-dependent growth rate changes The thicknesses of the (100)-, (110)-, and (111)-oriented singlecrystal SRO thin films were systematically controlled by modifying the number of pulsed laser shots used for the ablation of the target. Fig. 1(a)-(c) show the XRD θ-2θ patterns confirming the successful growth of the SRO epitaxial thin films, with the sharp diffraction peaks and clear Pendellӧsung fringes. The thickness of each film was derived based on the fringes and X-ray reflectivity (XRR) analyses complemented by the Pendellӧsung fringes and scanning electron microscopy (SEM). Fig. S1(a) shows the magnified XRR data for exemplary SRO thin films with different orientation. We also used SEM images to double-check the thicknesses as shown in Figs. S1(b). We obtained ~5% statistical and instrumental error in the thickness, from the various measurements. The total film thickness increased with the number of laser shots, evidenced by the Pendellӧsung fringes becoming denser. The analysis of the samples grown at the same time shows that the gap between the neighboring fringes is wider for the (111) orientation (Fig. 1(c)) than for the (100) orientation (Fig. 1(a)), indicating that the (100)-orientated SRO thin film is thicker although the same number of laser shots was employed. The surface-orientation-dependent epitaxial thin film growth rate is summarized in Fig. 1(d), as a function of the number of laser shots. A clear trend is observed as the crystallographic orientation changes from (100) to (110) and (111), while maintaining the linear relationship between the total thickness of the thin film and number of laser shots. The (100) orientation leads to the highest growth rate of 0.0038 nm/ shot, whereas the (110) and (111) orientations lead to lower values of 0.0034 and 0.0029 nm/shot, respectively. The crystallographic orientation dependence can also be synopsized by plotting the thickness difference between the (100)- and (111)-oriented SRO thin films (t(100) − t(111)), as shown in the inset of Fig. 1(d). The thickness difference evolves linearly with the number of laser shots.
2.2. Surface characterization using atomic force microscopy AFM measurements were performed using a commercial system (Park Systems, NX10) to examine the surface topography. To obtain surface potential images, KPFM measurements were performed using an AC modulation voltage of 2 Vrms at 17 kHz in the lift mode with a distance of 20 nm between the tip and the sample. In the KPFM measurements [16], the conductive Pt/Cr coated tip (BudgetSensors, Multi75E-G) was used. 2.3. Surface energy calculation To calculate the surface energies of the (100), (110), and (111) facets of SRO, we performed spin-polarized DFT calculations with the VASP code [17,18] and Perdew-Burke-Ernzerhof exchange correlation functional [19]. The interaction between the ionic core and the valence electrons was expressed by the projector augmented wave method [20]. The valence electron functions were expanded with a plane wave basis up to an energy cut-off of 400 eV. The Brillouin zone was sampled at the Γ-point. The convergence criteria for the electronic structure and geometry were set to 10−4 eV and 0.05 eV/Å, respectively. We used the Gaussian smearing method with a finite temperature width of 0.05 eV to improve the convergence of the states near the Fermi level. The (100) surface of SRO was modeled with a flat (3 × 3) slab with 9 atomic layers, while the (110) and (111) surfaces were modeled with 2
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Fig. 1. Thickness analysis of the (a) (100)-, (b) (110)-, and (c) (111)-oriented SRO thin films grown by PLE. The XRD and XRR results indicate a gradual increase in the film thickness with the number of laser shots. (d) Surface-orientation-dependent growth rates for the (100) (black), (110) (red), and (111) (blue) orientations. The inset shows the thickness difference between the (100) and (111) orientations. The dashed lines indicate linear fit of the thickness as a function of the number of laser shots.
interaction strength leads to a three-dimensional (3D) island (VolmerWeber) growth. The intermediate joint island (Stranski-Krastanove) growth is realized for the intermediate case. The competition between these interaction strengths can be modulated by the substrate temperature, energy of the incoming adatoms, width of the atomic step of the substrate, etc. [29–31]. In our case, these different growth modes were observed among the SRO epitaxial thin films with (100), (110), and (111) orientations. Fig. 2 presents AFM topography images (left panels) of the SRO thin film surfaces with different crystallographic orientations, showing the growth mode variation. The image in Fig. 2(a) for the (100) orientation clearly shows a step-flow growth, a type of 2D growth (Frank-van der Merwe), typically observed for the SRO thin film grown on a (100) STO substrate [32,33]. On the other hand, the image in Fig. 2(c) for the (111) orientation indicates a 3D island growth (Volmer-Weber), with the absence of the step-and-terrace structure of the substrate, but with small islands formed on the surface. Several previous studies also demonstrated 3D island growths of (relatively thick) SRO films with the (111) orientation [8,10,34], consistent with our result. The topography in Fig. 2(b) for the (110) orientation is in between the two modes discussed previously, suggesting the Stranski-Krastanov growth. The result can also be implied from the varying adatom-surface interaction through the polarity difference of each orientation, i.e., (SrO)0/(RuO2)0 for (100), (SrRuO)4+/(O2)4− for (110), and (Ru)4+/(SrO3) 4- for (111) orientation, respectively. The non-polar state in (100) results in a stable 2D growth while the polar states in (110) and (111) result in 3D growth as imagined. We note that the step width of the substrate could also
The lower growth rates for the (110) and (111) orientations than those for the (100) orientation are rather unexpected as the growth was performed at the same time, manifesting the same growth parameters. It is possible that the distinctive epitaxial strain states owing to the different crystallographic-orientation-induced distortion can lead to the different growth rates of the thin films. However, considering the pseudocubic SRO thin film on the cubic STO substrate, the epitaxial strain is only ~0.64% [24]. An additional variation owing to the orientation-dependent distortion would be too small to explain the observed difference (> 20%) between the (100)- and (111)-oriented thin films. Moreover, the films become relaxed above the thickness of ~80 nm [25], yet the linear trend is maintained, indicating that factors other than the epitaxial strain should contribute to the change in the growth rate. 3.2. Surface-property-dependent growth characterization In addition to the growth rate, the modification of the surface orientation influenced the growth mode of the thin films. Classically, the thin film growths can be categorized into three modes: Frank-van der Merwe [26], Stranski-Krastanov [27], and Volmer-Weber [28]. These modes can be classified depending on the competition between the adatom-adatom and adatom-surface interaction strengths. For a relatively high adatom-surface interaction strength compared to the adatom-adatom interaction strength, a two-dimensional (2D) (Frankvan der Merwe) growth is achieved. In contrast, a relatively low adatom-surface interaction strength compared to the adatom-adatom 3
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Fig. 2. AFM topography and KPFM surface potential images of the (100)-, (110)-, and (111)-oriented SRO thin films. AFM topography images for the (a) (100), (b) (110), and (c) (111) orientations showing the different growth modes (left panels). The scale bar represents 1 μm. The KPFM images (right panels) indicate the surface potentials, as summarized in (d). To minimize the substrate-induced potential difference, we chose the thickest film set (with 25,000 laser shots). Table 1 DFT-calculated surface energies of the SRO surfaces, unit-cell surface energy, and unit-cell nucleation barrier energies for the different surface orientations. The average surface energy (γ) is derived from the distinctive surface terminations. S is the in-plane surface area of the unit cell with the corresponding orientation. The formation energy (ΔEbulk = −5.76 eV) is the energy needed to form one unit cell of SRO. The unit-cell nucleation barrier energy (ΔGu) is derived using the classical nucleation theory. Surface orientation Termination
(100) SrO
Surface energy (eV/nm2) Average surface energy γ (eV/nm2) Unit-cell surface area S (nm2) Unit-cell surface energy Sγ (eV) Unit-cell nucleation barrier energy ΔGu (eV)
6.43 6.20 0.16 0.99 −4.77
(110) SrRuO
RuO2
5.98
11.96 9.59 0.22 2.09 −3.67
O2
7.22
(111) Ru
10.24 9.75 0.27 2.61 −3.15
SrO3
9.26
more stable owing to the slightly higher surface potential value.
influence the growth mode [29,35], and confirmed that all STO substrates used in this study (with different orientations) had similar step widths. Fig. S2 shows the AFM topography images of SRO films grown on (100)-, (110)-, and (111)-oriented STO substrates with various thicknesses. The results confirm that the growth mode is not much affected by the thickness for the same surface orientation, at least when the film is thicker than ~10 nm. The surface potential was obtained from the corresponding KPFM measurements, which showed a systematic change with the surface orientation, as shown in the right panels of Fig. 2. The relative difference in the surface potential can tell us the difference in the charge states depending on the crystallographic orientation which is closely related to the adatom-surface interaction and adatom mobility in film growth system. The KPFM results show the increase in surface potential with the change in orientation from (100) to (110) and (111), as summarized in Fig. 2(d). The lowest surface potential of ~82 mV observed for the (100)-oriented SRO (Fig. 2(a)) enables the film growth with a relatively high adatom-surface interaction strength, leading to the 2D step-flow growth. In contrast, higher surface potentials of ~140 and ~150 mV are observed for the (110)- and (111)-oriented SRO (Fig. 2(b) and (c)), respectively. The relatively low adatom-surface interaction strength leads to congregation of the adatoms forming the islands. The minor potential difference of ~10 mV between the (110)and (111)-oriented samples leads to a similar 3D island growth, yet, it seems that the islands formed in the (111)-oriented SRO surface are
3.3. Nucleation barrier energy calculation We employ the classical nucleation theory to understand the surface-orientation-dependent growth of the SRO thin films. The nucleation governs not only the initial growth of the thin film but also the continuous stacking of the lattice, as the thin film continues to form on a newly formed surface as the growth proceeds. According to the classical nucleation theory [36,37], the film nucleation can be modeled in terms of the nucleation barrier energy (ΔGn) with a certain nucleus radius (r). ΔGn can be expressed as
4 ΔGn = − πr 3ΔGv + πr 2γ , 3
(3)
where ΔGv, r, and γ are the Gibbs free energy per unit volume, radius of the nucleus, and surface energy per unit area, respectively. Therefore, ΔGn is determined by the competition between the formation energy (first term) and interfacial energy (second term). The formation energy is dependent on the nucleus radius and Gibbs free energy of the film material. The interfacial energy is also dependent on the size of the nucleus and surface energy of the film material. Based on the energetic competition, we can qualitatively understand the experimentally observed change in the thin film growth. First, the growth is favored when ΔGn is small. This is the case when the interfacial energy term is small. 4
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the unit-cell nucleation barrier energy of the SRO thin film derived from ΔGn for the macroscopic nucleation (Eq. (3)). The first term (formation energy) of Eq. (4), ΔEbulk represents the formation energy of a single bulk SRO unit cell from separated SRO clusters. The second term of Eq. (4), Sγ, represents the surface energy of one SRO unit cell, determined by the surface energy and unit cell in-plane area through the surfaceorientation-dependent atomic and chemical structures (Table 1). The surface energies of the (100)-, (110)-, and (111)-oriented SRO structures were calculated by spin-polarized DFT calculations considering the distinctive surface atomic terminations (SrO and RuO2 for (100), SrRuO and O2 for (110), and Ru and SrO3 for (111)). The surface energies are 6.43 and 5.98 (Esurface(100)−SrO and Esurface(100)−RuO2 for the (100) orientation), 11.96 and 7.22 (Esurface(110)−SrRuO and Esurface(110)−O2 for the (110) orientation), and 10.24 and 9.26 (Esurface(111)−Ru and Esurface(111)−SrO3 for the (111) orientation) eV/nm2, respectively (Table 1). As anticipated, the non-polar (100) facets had lower surface energy compared to those of (110) and (111) facets. As the exact surface atomic termination is not known except for the (100) orientation, the average surface energies (γ) of 6.20, 9.59, and 9.75 eV/ nm2 are presented for the (100), (110), and (111) orientations, respectively. S depends on the unit cell in-plane geometries of square, rectangular, and triangular shapes, yielding values of 0.16, 0.22, and 0.27 nm2 for the (100), (110), and (111) orientations, respectively. Considering the distinct S and γ values, the unit-cell surface energy, Sγ, of SRO varies with the crystallographic surface orientation (0.99, 2.10, and 2.61 eV); the resultant ΔGu values are −4.77, −3.67, and − 3.15 eV for the (100), (110), and (111) orientations, respectively. The systematic increase in ΔGu with the surface orientation change from (100) to (110) to (111) reflects the experimentally observed differences in the growth rate and growth mode of the SRO thin film. In particular, the orientation-dependent change in the unit-cell surface energy governs the characteristic growth on the atomistic scale. The lowest Sγ of the (100) orientation leads to the lowest ΔGu, corresponding to a high growth rate (Fig. 1). The relatively small Sγ of the (100) orientation, compared to those of the (110) and (111) orientations, further suggests a higher adatom-surface interaction strength, leading to a 2D growth of the SRO thin film, as shown in Fig. 2. As other important growth parameters, e.g., the temperature, pressure, incoming atomic concentration, and energy, are not explicitly included in the calculation, quantitative determination of the theoretical values would be rather insignificant. Nevertheless, the relative unit-cell surface energies among the (100), (110), and (111) orientations provide consistent understanding of the orientation-dependent growth mode of the SRO epitaxial thin film.
Fig. 3. Topography and crystallographic surface orientation information images of the epitaxial polycrystalline SRO thin film. (a) EBSD inverse pole figure and (b) AFM topography images of the same region of the polycrystalline SRO thin film. The black square region in (a) was chosen to compare the grains with the surface orientations close to the three representative surface orientations of (100) (red), (110) (green), and (111) (blue). (c) Profile along the black line in (b) covering the red ((100)-like), green ((110)-like), and blue ((111)like) grains. It shows the thickness variations of ~30 to ~50 nm depending on the orientation, similar to the variations of the single-crystal thin films.
3.4. Combinatorial investigation using polycrystalline epitaxial SRO thin films Our analysis was further tested combinatorially, taking advantage of the polycrystalline epitaxial SRO thin films [9]. Fig. 3 shows a colored inverse pole figure map obtained by EBSD (Fig. 3(a)) and corresponding AFM topography image (Fig. 3(b)) of the polycrystalline epitaxial SRO thin film. The results enable us to combinatorially compare the thickness (growth rate) and surface morphology of the single-crystalline epitaxial grains, depending on the various surface crystallographic orientations. For example, using the region marked by the black square in Fig. 3(a), we can simultaneously compare the thicknesses of the grains oriented close to the (100) (red), (110) (green), and (111) (blue) orientations. The profile along the black line in Fig. 3(b) is shown in Fig. 3(c). A clear difference in the growth rate is observed, consistent with the result obtained for the single-crystal thin films. The redmarked grain ((100)-like) has the highest, while the blue-marked grain ((111)-like) has the lowest growth rate. The difference in the surface morphology can also be inferred from the polycrystalline epitaxial SRO thin film. The advantage of the polycrystalline epitaxial SRO thin film is
When the interfacial energy is increased, more energy should be supplied for the growth, which implies a slower growth. On the other hand, a relatively small formation energy compared to the interfacial energy would mean a larger adatom-adatom interaction strength than the adatom-surface interaction strength, yielding an island growth. For application of the classical nucleation theory for the atomicscale PLE growth, we further consider the following modifications to the equation, while the qualitative principle is identical. First, ΔGn in Eq. (3) can be modified to the unit-cell nucleation barrier energy,
ΔGu = ∆Ebulk + Sγ ,
(4)
where ΔEbulk and S are the formation energy of the unit cell and unit cell in-plane area, respectively. Using Eq. (4), we can calculate ΔGu for 5
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the ability to investigate the crystallographic orientations beyond the lowest-Miller-index surfaces, which could be analyzed in a following study. Nevertheless, the qualitative tendency is clear. With the change in the surface orientation from (100) to (111), the thin film growth rate decreases and the 2D growth approaches to a 3D island growth.
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Hesse, Substrate temperature dependence of structure and resistivity of SrRuO3 thin films grown by pulsed laser deposition on (100) SrTiO3, J. Mater. Res. 14 (2011) 4385–4394. [31] V.S. Vidhya, V. Malathy, T. Balasubramanian, V. Saaminathan, C. Sanjeeviraja, M. Jayachandran, Influence of RF power on the growth mechanism, preferential orientation and optoelectronic properties of nanocrystalline ITO films, Curr. Appl. Phys. 11 (2011) 286–294. [32] J. Choi, C.B. Eom, G. Rijnders, H. Rogalla, D.H.A. Blank, Growth mode transition from layer by layer to step flow during the growth of heteroepitaxial SrRuO3 on (001) SrTiO3, Appl. Phys. Lett. 79 (2001) 1447–1449. [33] B. Kim, S.A. Lee, D. Seol, W.S. Choi, Y. Kim, Ex-situ atomic force microscopy on the growth mode of SrRuO3 epitaxial thin film, Curr. Appl. Phys. 17 (2017) 1721–1726. [34] D. Rubi, A.H.G. Vlooswijk, B. 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4. Conclusion We investigated the crystallographic surface-orientation-dependent growth rate and growth modes of the SRO epitaxial thin films on the STO substrates using PLE. The growth characteristics were analyzed using XRD, AFM, and KPFM for the representative crystallographic surface orientations ((100), (110), and (111)) of the single-crystal thin films. With the change in the orientation from (100) to (110) to (111), the growth rate significantly decreased and the growth mode evolved from the 2D to the 3D growth. According to the first-principles calculation and atomistically modified nucleation barrier energy calculation, the unit-cell surface energy of SRO increased with the change in the orientation from (100) to (110) to (111), consistent with the experimentally observed changes. The systematic understanding of the singlecrystal thin films was further applied to the epitaxial polycrystalline thin film with arbitrarily oriented grains. Our study provides a design principle for functional oxide epitaxial thin films and heterostructures with different surface orientations necessary for conventional and future oxide opto-electronic applications. Acknowledgments This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF2019R1A2B5B02004546 and NRF-2019R1A6A1A03033215) and Programme Asie 127096 supported by the Conseil Régional de Normandie, and PHC Star 36453VL (AD & WP). Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.apsusc.2019.143924. References [1] G. Koster, L. Klein, W. Siemons, G. Rijnders, J.S. Dodge, C.-B. Eom, D.H.A. Blank, M.R. Beasley, Structure, physical properties, and applications of SrRuO3 thin films, Rev. Mod. Phys. 84 (2012) 253–298. [2] S.A. Lee, S. Oh, J.-Y. Hwang, M. Choi, C. Youn, J.W. Kim, S.H. Chang, S. Woo, J.S. Bae, S. Park, Y.-M. Kim, S. Lee, T. Choi, S.W. Kim, W.S. Choi, Enhanced electrocatalytic activity via phase transitions in strongly correlated SrRuO3 thin films, Energy Environ. Sci. 10 (2017) 924–930. [3] S.A. Lee, S. Oh, J. Lee, J.-Y. Hwang, J. Kim, S. Park, J.-S. Bae, T.E. Hong, S. Lee, S.W. Kim, W.N. Kang, W.S. Choi, Tuning electromagnetic properties of SrRuO3 epitaxial thin films via atomic control of cation vacancies, Sci. Rep. 7 (2017) 11583. [4] Y.J. Chang, C.H. Kim, S.H. Phark, Y.S. Kim, J. Yu, T.W. Noh, Fundamental thickness limit of itinerant ferromagnetic SrRuO3 thin films, Phys. Rev. Lett. 103 (2009) 057201. [5] D. Kan, Y. Shimakawa, Strain effect on structural transition in SrRuO3 epitaxial thin films, Cryst. Growth Des. 11 (2011) 5483–5487. [6] K. Ishigami, K. Yoshimatsu, D. Toyota, M. Takizawa, T. Yoshida, G. Shibata, T. Harano, Y. Takahashi, T. Kadono, V.K. Verma, V.R. Singh, Y. Takeda, T. Okane, Y. Saitoh, H. Yamagami, T. Koide, M. Oshima, H. Kumigashira, A. Fujimori, Thickness-dependent magnetic properties and strain-induced orbital magnetic moment in SrRuO3 thin films, Phys. Rev. B 92 (2015) 064402. [7] P. Kaur, K.K. Sharma, R. Pandit, R.J. Choudhary, R. Kumar, Structural, electrical, and magnetic properties of SrRuO3 thin films, Appl. Phys. Lett. 104 (2014) 081608.
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Surface-orientation-dependent growth of SrRuO3 epitaxial thin films a
b
c,d
a
e
T e
Sungmin Woo , Hyuk Choi , Seunghun Kang , Jegon Lee , Adrian David , Wilfrid Prellier , ⁎ Yunseok Kimc,d, Hyun You Kimb, Woo Seok Choia, a
Department of Physics, Sungkyunkwan University, Suwon 16419, Republic of Korea Department of Material Science and Engineering, Chungnam National University, Daejeon 34134, Republic of Korea School of Advanced Materials Science and Engineering, Sungkyunkwan University, Suwon 16419, Republic of Korea d Research Center for Advanced Materials Technology, Sungkyunkwan University, Suwon 16419, Republic of Korea e Laboratorie CRISMAT, CNRS UMR 6508, ENSICAEN, Normandie Universite, 6 Bd Marechal Juin, F-14050 Caen Cedex 4, France b c
A R T I C LE I N FO
A B S T R A C T
Keywords: Pulsed laser epitaxy Crystallographic orientation Surface energy Growth kinetics SrRuO3
The growth of epitaxial transition metal oxide thin films depends on various parameters including the substrate temperature, oxygen partial pressure, kinetics of incoming adatoms, Gibbs free energy, and surface energy. Naturally, the change in the crystallographic surface orientation with a distinctive surface energy also influences the growth rate and growth mode of the epitaxial thin films substantially. Using perovskite SrRuO3 as a model system, we studied the growth characteristics by changing the surface orientation of the SrTiO3 substrate. Employing X-ray diffraction and surface atomic and Kelvin probe force microscopy (KPFM), we observed a systematic decrease in the growth rate and a modification in the growth mode from a two-dimensional growth to a three-dimensional island growth with the change in the surface orientation from (100) to (110) to (111). A spin-polarized density functional theory calculation demonstrated the corresponding difference in the surface energy, which was also confirmed experimentally by the KPFM measurements. The difference in the surface energy could explain the observed change in the growth kinetics, based on the modified classical nucleation theory. A combinatorial method using a polycrystalline epitaxial thin film was employed to generalize our understanding of the crystallographic surface-orientation-dependent thin film growth.
1. Introduction Perovskite SrRuO3 (SRO) is a strongly correlated itinerant ferromagnet, widely employed as a metallic electrode for oxide electronics. Bulk SRO has an orthorhombic crystal structure with ferromagnetism arising below the Curie temperature (TC) of ~160 K [1]. When epitaxially grown on different perovskite substrates, it exhibits structural transition from an orthorhombic to a tetragonal structure as a function of the epitaxial strain and thickness, which influences its electronic and magnetic properties [2–6]. The crystallographic surface orientation of the epitaxial SRO thin film is also crucial for its physical behaviors. A distinctive structural distortion can be anticipated, i.e., tetragonal distortion for (100)-, monoclinic distortion for (110)-, and trigonal distortion for (111)-oriented epitaxial thin films. The difference in the structural distortion affects the electronic structure and magnetic exchange coupling, eventually modifying the itinerant nature and ferromagnetic behavior [7–10]. Therefore, precise understanding and control of the growth of the epitaxial thin films with various orientations are highly necessary.
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Indeed, the atomic-scale control of the perovskite oxide thin film growth is of prime importance for the state-of-the-art heterostructure fabrication techniques. The growth rate of the perovskite oxide thin film is determined by the combination of several parameters including the surface energies of the substrate and thin film, growth temperature and pressure, incoming atomic concentration and energy, Gibbs free energy, and nucleation and desorption energies of the atoms. The orientation dependence of the growth rate has been studied, for growth methods such as chemical vapor deposition and sputtering using substrates with different surface orientations [11–14]. However, most of the films used to analyze the growth mechanism were on the micrometer scale, such that the atomic-scale information was not accessible. In addition, the growth mode variation with the crystallographic orientation has not yet been explained in detail for a systematically controlled growth system. For better understanding of the orientationdependent growth mechanism leading to fine synthetic crystals with atomic-scale precision, studies exploiting the well-established pulsed laser epitaxy (PLE) are required. In this study, we investigated the relationship between the surface
Corresponding author. E-mail address: [email protected] (W.S. Choi).
https://doi.org/10.1016/j.apsusc.2019.143924 Received 14 June 2019; Received in revised form 28 August 2019; Accepted 9 September 2019 Available online 10 September 2019 0169-4332/ © 2019 Elsevier B.V. All rights reserved.
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energy and growth mode (and rate) of SRO epitaxial thin films by employing various surface orientations, with a specific focus on the atomic-scale precision epitaxy. We fabricated epitaxial SRO thin films on SrTiO3 (STO) substrates with (100), (110), and (111) surface orientations using PLE. Various film thicknesses in the range of a few to a hundred nanometers were investigated. Using atomic force microscopy (AFM) and Kelvin probe force microscopy (KPFM), we confirmed the growth mode variation with the surface potential of each orientation. The surface energies of the (100)-, (110)-, and (111)-oriented SRO films were calculated by a spin-polarized density functional theory (DFT) calculation, yielding results consistent with the experimentally obtained values. Based on the modified classical nucleation theory, the different unit-cell surface energies provided comprehensive understanding of the orientation-dependent growth mechanism. Furthermore, we employed epitaxial polycrystalline SRO thin films for combinatorial confirmation of the crystallographic surface-orientationdependent growth.
(3 × 3) slabs with 12 atomic layers. The topmost and bottommost layers of the model slabs were released in the optimization to calculate the energy associated with the surface relaxation. As a stoichiometric SRO slab cannot be sliced into two SRO slabs with two identical surface morphologies, we applied an average estimation method, which was also applied in our previous study [15]. Two different surface terminations were considered for each surface: RuO2 or SrO termination for (100), SrRuO or O2 termination for (110), and Ru or SrO3 termination for (111). As both surface terminations simultaneously form upon cleavage of a stoichiometric SRO slab, we initially calculated the cleavage energy of SRO, Ecleavage, and equally distributed it to both terminations. For example, the unrelaxed internal cleavage energy of the (100) surface was calculated as [21–23]. (100) Ecleavage =
1 (100) − RuO2 (100) − SrO [E + Eslab − x ΔEbulk ] 4 slab (100)−RuO2
(1)
(100)−SrO
and Eslab are the energies of the unrelaxed where Eslab RuO2- and SrO-terminated surfaces, respectively, ΔEbulk represents the formation energy of a single bulk SRO unit cell from separated SRO clusters obtained from the calculation (ΔEbulk = −5.76 eV), and x is the number of total unit cells in both slabs. For the (100), (110), and (111) surfaces, the x values were 81, 81, and 99, respectively. The factor of 1/ 4 is included as four surfaces were created upon cleavage. Based on Ecleavage and ΔEbulk, for example, the surfaces energy of the SrO-terminated (100) surface was calculated as
2. Methods 2.1. Epitaxial thin film growth and structural characterization High-quality epitaxial single-crystal and polycrystalline SRO thin films were synthesized on STO single-crystal ((100), (110), and (111)) and polycrystalline substrates using PLE at 700 °C, respectively. The substrates were treated to obtain atomically flat surfaces [15]. An excimer laser (248 nm; IPEX 864, Lightmachinery) with a fluence of 1.5 J/ cm2 and repetition rate of 5 Hz was used. The thin films were grown at an oxygen partial pressure of 100 mTorr, which results in a stoichiometric SRO thin film [3]. Each set of thin films (with the same number of laser shots) was fabricated at the same time. The atomic structures, crystal orientations, epitaxy relationships, and thicknesses of the SRO thin films were characterized using high-resolution X-ray diffraction (XRD, Malvern PANalytical, X'Pert Pro MRD XL) and electron backscattering diffraction (EBSD, JSM7000F). For the EBSD, the sample was mounted at an angle of 70° from the surface normal within a scanning electron microscope operated at 20 kV. The probe current of the aperture was 1 × 10−8 A. The Kikuchi diagrams were recorded with a beam step size of 0.5 μm. A commercial program (OIM Analysis 5.31) was used for image processing.
(100) − SrO (100) Esurface = Ecleavage +
1 (100) − SrO ∆Erelaxation 2
(2)
where ΔErelaxation(100)−SrO is the energy released upon the relaxation of the slab. 3. Results and discussion 3.1. Crystallographic surface-orientation-dependent growth rate changes The thicknesses of the (100)-, (110)-, and (111)-oriented singlecrystal SRO thin films were systematically controlled by modifying the number of pulsed laser shots used for the ablation of the target. Fig. 1(a)-(c) show the XRD θ-2θ patterns confirming the successful growth of the SRO epitaxial thin films, with the sharp diffraction peaks and clear Pendellӧsung fringes. The thickness of each film was derived based on the fringes and X-ray reflectivity (XRR) analyses complemented by the Pendellӧsung fringes and scanning electron microscopy (SEM). Fig. S1(a) shows the magnified XRR data for exemplary SRO thin films with different orientation. We also used SEM images to double-check the thicknesses as shown in Figs. S1(b). We obtained ~5% statistical and instrumental error in the thickness, from the various measurements. The total film thickness increased with the number of laser shots, evidenced by the Pendellӧsung fringes becoming denser. The analysis of the samples grown at the same time shows that the gap between the neighboring fringes is wider for the (111) orientation (Fig. 1(c)) than for the (100) orientation (Fig. 1(a)), indicating that the (100)-orientated SRO thin film is thicker although the same number of laser shots was employed. The surface-orientation-dependent epitaxial thin film growth rate is summarized in Fig. 1(d), as a function of the number of laser shots. A clear trend is observed as the crystallographic orientation changes from (100) to (110) and (111), while maintaining the linear relationship between the total thickness of the thin film and number of laser shots. The (100) orientation leads to the highest growth rate of 0.0038 nm/ shot, whereas the (110) and (111) orientations lead to lower values of 0.0034 and 0.0029 nm/shot, respectively. The crystallographic orientation dependence can also be synopsized by plotting the thickness difference between the (100)- and (111)-oriented SRO thin films (t(100) − t(111)), as shown in the inset of Fig. 1(d). The thickness difference evolves linearly with the number of laser shots.
2.2. Surface characterization using atomic force microscopy AFM measurements were performed using a commercial system (Park Systems, NX10) to examine the surface topography. To obtain surface potential images, KPFM measurements were performed using an AC modulation voltage of 2 Vrms at 17 kHz in the lift mode with a distance of 20 nm between the tip and the sample. In the KPFM measurements [16], the conductive Pt/Cr coated tip (BudgetSensors, Multi75E-G) was used. 2.3. Surface energy calculation To calculate the surface energies of the (100), (110), and (111) facets of SRO, we performed spin-polarized DFT calculations with the VASP code [17,18] and Perdew-Burke-Ernzerhof exchange correlation functional [19]. The interaction between the ionic core and the valence electrons was expressed by the projector augmented wave method [20]. The valence electron functions were expanded with a plane wave basis up to an energy cut-off of 400 eV. The Brillouin zone was sampled at the Γ-point. The convergence criteria for the electronic structure and geometry were set to 10−4 eV and 0.05 eV/Å, respectively. We used the Gaussian smearing method with a finite temperature width of 0.05 eV to improve the convergence of the states near the Fermi level. The (100) surface of SRO was modeled with a flat (3 × 3) slab with 9 atomic layers, while the (110) and (111) surfaces were modeled with 2
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Fig. 1. Thickness analysis of the (a) (100)-, (b) (110)-, and (c) (111)-oriented SRO thin films grown by PLE. The XRD and XRR results indicate a gradual increase in the film thickness with the number of laser shots. (d) Surface-orientation-dependent growth rates for the (100) (black), (110) (red), and (111) (blue) orientations. The inset shows the thickness difference between the (100) and (111) orientations. The dashed lines indicate linear fit of the thickness as a function of the number of laser shots.
interaction strength leads to a three-dimensional (3D) island (VolmerWeber) growth. The intermediate joint island (Stranski-Krastanove) growth is realized for the intermediate case. The competition between these interaction strengths can be modulated by the substrate temperature, energy of the incoming adatoms, width of the atomic step of the substrate, etc. [29–31]. In our case, these different growth modes were observed among the SRO epitaxial thin films with (100), (110), and (111) orientations. Fig. 2 presents AFM topography images (left panels) of the SRO thin film surfaces with different crystallographic orientations, showing the growth mode variation. The image in Fig. 2(a) for the (100) orientation clearly shows a step-flow growth, a type of 2D growth (Frank-van der Merwe), typically observed for the SRO thin film grown on a (100) STO substrate [32,33]. On the other hand, the image in Fig. 2(c) for the (111) orientation indicates a 3D island growth (Volmer-Weber), with the absence of the step-and-terrace structure of the substrate, but with small islands formed on the surface. Several previous studies also demonstrated 3D island growths of (relatively thick) SRO films with the (111) orientation [8,10,34], consistent with our result. The topography in Fig. 2(b) for the (110) orientation is in between the two modes discussed previously, suggesting the Stranski-Krastanov growth. The result can also be implied from the varying adatom-surface interaction through the polarity difference of each orientation, i.e., (SrO)0/(RuO2)0 for (100), (SrRuO)4+/(O2)4− for (110), and (Ru)4+/(SrO3) 4- for (111) orientation, respectively. The non-polar state in (100) results in a stable 2D growth while the polar states in (110) and (111) result in 3D growth as imagined. We note that the step width of the substrate could also
The lower growth rates for the (110) and (111) orientations than those for the (100) orientation are rather unexpected as the growth was performed at the same time, manifesting the same growth parameters. It is possible that the distinctive epitaxial strain states owing to the different crystallographic-orientation-induced distortion can lead to the different growth rates of the thin films. However, considering the pseudocubic SRO thin film on the cubic STO substrate, the epitaxial strain is only ~0.64% [24]. An additional variation owing to the orientation-dependent distortion would be too small to explain the observed difference (> 20%) between the (100)- and (111)-oriented thin films. Moreover, the films become relaxed above the thickness of ~80 nm [25], yet the linear trend is maintained, indicating that factors other than the epitaxial strain should contribute to the change in the growth rate. 3.2. Surface-property-dependent growth characterization In addition to the growth rate, the modification of the surface orientation influenced the growth mode of the thin films. Classically, the thin film growths can be categorized into three modes: Frank-van der Merwe [26], Stranski-Krastanov [27], and Volmer-Weber [28]. These modes can be classified depending on the competition between the adatom-adatom and adatom-surface interaction strengths. For a relatively high adatom-surface interaction strength compared to the adatom-adatom interaction strength, a two-dimensional (2D) (Frankvan der Merwe) growth is achieved. In contrast, a relatively low adatom-surface interaction strength compared to the adatom-adatom 3
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Fig. 2. AFM topography and KPFM surface potential images of the (100)-, (110)-, and (111)-oriented SRO thin films. AFM topography images for the (a) (100), (b) (110), and (c) (111) orientations showing the different growth modes (left panels). The scale bar represents 1 μm. The KPFM images (right panels) indicate the surface potentials, as summarized in (d). To minimize the substrate-induced potential difference, we chose the thickest film set (with 25,000 laser shots). Table 1 DFT-calculated surface energies of the SRO surfaces, unit-cell surface energy, and unit-cell nucleation barrier energies for the different surface orientations. The average surface energy (γ) is derived from the distinctive surface terminations. S is the in-plane surface area of the unit cell with the corresponding orientation. The formation energy (ΔEbulk = −5.76 eV) is the energy needed to form one unit cell of SRO. The unit-cell nucleation barrier energy (ΔGu) is derived using the classical nucleation theory. Surface orientation Termination
(100) SrO
Surface energy (eV/nm2) Average surface energy γ (eV/nm2) Unit-cell surface area S (nm2) Unit-cell surface energy Sγ (eV) Unit-cell nucleation barrier energy ΔGu (eV)
6.43 6.20 0.16 0.99 −4.77
(110) SrRuO
RuO2
5.98
11.96 9.59 0.22 2.09 −3.67
O2
7.22
(111) Ru
10.24 9.75 0.27 2.61 −3.15
SrO3
9.26
more stable owing to the slightly higher surface potential value.
influence the growth mode [29,35], and confirmed that all STO substrates used in this study (with different orientations) had similar step widths. Fig. S2 shows the AFM topography images of SRO films grown on (100)-, (110)-, and (111)-oriented STO substrates with various thicknesses. The results confirm that the growth mode is not much affected by the thickness for the same surface orientation, at least when the film is thicker than ~10 nm. The surface potential was obtained from the corresponding KPFM measurements, which showed a systematic change with the surface orientation, as shown in the right panels of Fig. 2. The relative difference in the surface potential can tell us the difference in the charge states depending on the crystallographic orientation which is closely related to the adatom-surface interaction and adatom mobility in film growth system. The KPFM results show the increase in surface potential with the change in orientation from (100) to (110) and (111), as summarized in Fig. 2(d). The lowest surface potential of ~82 mV observed for the (100)-oriented SRO (Fig. 2(a)) enables the film growth with a relatively high adatom-surface interaction strength, leading to the 2D step-flow growth. In contrast, higher surface potentials of ~140 and ~150 mV are observed for the (110)- and (111)-oriented SRO (Fig. 2(b) and (c)), respectively. The relatively low adatom-surface interaction strength leads to congregation of the adatoms forming the islands. The minor potential difference of ~10 mV between the (110)and (111)-oriented samples leads to a similar 3D island growth, yet, it seems that the islands formed in the (111)-oriented SRO surface are
3.3. Nucleation barrier energy calculation We employ the classical nucleation theory to understand the surface-orientation-dependent growth of the SRO thin films. The nucleation governs not only the initial growth of the thin film but also the continuous stacking of the lattice, as the thin film continues to form on a newly formed surface as the growth proceeds. According to the classical nucleation theory [36,37], the film nucleation can be modeled in terms of the nucleation barrier energy (ΔGn) with a certain nucleus radius (r). ΔGn can be expressed as
4 ΔGn = − πr 3ΔGv + πr 2γ , 3
(3)
where ΔGv, r, and γ are the Gibbs free energy per unit volume, radius of the nucleus, and surface energy per unit area, respectively. Therefore, ΔGn is determined by the competition between the formation energy (first term) and interfacial energy (second term). The formation energy is dependent on the nucleus radius and Gibbs free energy of the film material. The interfacial energy is also dependent on the size of the nucleus and surface energy of the film material. Based on the energetic competition, we can qualitatively understand the experimentally observed change in the thin film growth. First, the growth is favored when ΔGn is small. This is the case when the interfacial energy term is small. 4
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the unit-cell nucleation barrier energy of the SRO thin film derived from ΔGn for the macroscopic nucleation (Eq. (3)). The first term (formation energy) of Eq. (4), ΔEbulk represents the formation energy of a single bulk SRO unit cell from separated SRO clusters. The second term of Eq. (4), Sγ, represents the surface energy of one SRO unit cell, determined by the surface energy and unit cell in-plane area through the surfaceorientation-dependent atomic and chemical structures (Table 1). The surface energies of the (100)-, (110)-, and (111)-oriented SRO structures were calculated by spin-polarized DFT calculations considering the distinctive surface atomic terminations (SrO and RuO2 for (100), SrRuO and O2 for (110), and Ru and SrO3 for (111)). The surface energies are 6.43 and 5.98 (Esurface(100)−SrO and Esurface(100)−RuO2 for the (100) orientation), 11.96 and 7.22 (Esurface(110)−SrRuO and Esurface(110)−O2 for the (110) orientation), and 10.24 and 9.26 (Esurface(111)−Ru and Esurface(111)−SrO3 for the (111) orientation) eV/nm2, respectively (Table 1). As anticipated, the non-polar (100) facets had lower surface energy compared to those of (110) and (111) facets. As the exact surface atomic termination is not known except for the (100) orientation, the average surface energies (γ) of 6.20, 9.59, and 9.75 eV/ nm2 are presented for the (100), (110), and (111) orientations, respectively. S depends on the unit cell in-plane geometries of square, rectangular, and triangular shapes, yielding values of 0.16, 0.22, and 0.27 nm2 for the (100), (110), and (111) orientations, respectively. Considering the distinct S and γ values, the unit-cell surface energy, Sγ, of SRO varies with the crystallographic surface orientation (0.99, 2.10, and 2.61 eV); the resultant ΔGu values are −4.77, −3.67, and − 3.15 eV for the (100), (110), and (111) orientations, respectively. The systematic increase in ΔGu with the surface orientation change from (100) to (110) to (111) reflects the experimentally observed differences in the growth rate and growth mode of the SRO thin film. In particular, the orientation-dependent change in the unit-cell surface energy governs the characteristic growth on the atomistic scale. The lowest Sγ of the (100) orientation leads to the lowest ΔGu, corresponding to a high growth rate (Fig. 1). The relatively small Sγ of the (100) orientation, compared to those of the (110) and (111) orientations, further suggests a higher adatom-surface interaction strength, leading to a 2D growth of the SRO thin film, as shown in Fig. 2. As other important growth parameters, e.g., the temperature, pressure, incoming atomic concentration, and energy, are not explicitly included in the calculation, quantitative determination of the theoretical values would be rather insignificant. Nevertheless, the relative unit-cell surface energies among the (100), (110), and (111) orientations provide consistent understanding of the orientation-dependent growth mode of the SRO epitaxial thin film.
Fig. 3. Topography and crystallographic surface orientation information images of the epitaxial polycrystalline SRO thin film. (a) EBSD inverse pole figure and (b) AFM topography images of the same region of the polycrystalline SRO thin film. The black square region in (a) was chosen to compare the grains with the surface orientations close to the three representative surface orientations of (100) (red), (110) (green), and (111) (blue). (c) Profile along the black line in (b) covering the red ((100)-like), green ((110)-like), and blue ((111)like) grains. It shows the thickness variations of ~30 to ~50 nm depending on the orientation, similar to the variations of the single-crystal thin films.
3.4. Combinatorial investigation using polycrystalline epitaxial SRO thin films Our analysis was further tested combinatorially, taking advantage of the polycrystalline epitaxial SRO thin films [9]. Fig. 3 shows a colored inverse pole figure map obtained by EBSD (Fig. 3(a)) and corresponding AFM topography image (Fig. 3(b)) of the polycrystalline epitaxial SRO thin film. The results enable us to combinatorially compare the thickness (growth rate) and surface morphology of the single-crystalline epitaxial grains, depending on the various surface crystallographic orientations. For example, using the region marked by the black square in Fig. 3(a), we can simultaneously compare the thicknesses of the grains oriented close to the (100) (red), (110) (green), and (111) (blue) orientations. The profile along the black line in Fig. 3(b) is shown in Fig. 3(c). A clear difference in the growth rate is observed, consistent with the result obtained for the single-crystal thin films. The redmarked grain ((100)-like) has the highest, while the blue-marked grain ((111)-like) has the lowest growth rate. The difference in the surface morphology can also be inferred from the polycrystalline epitaxial SRO thin film. The advantage of the polycrystalline epitaxial SRO thin film is
When the interfacial energy is increased, more energy should be supplied for the growth, which implies a slower growth. On the other hand, a relatively small formation energy compared to the interfacial energy would mean a larger adatom-adatom interaction strength than the adatom-surface interaction strength, yielding an island growth. For application of the classical nucleation theory for the atomicscale PLE growth, we further consider the following modifications to the equation, while the qualitative principle is identical. First, ΔGn in Eq. (3) can be modified to the unit-cell nucleation barrier energy,
ΔGu = ∆Ebulk + Sγ ,
(4)
where ΔEbulk and S are the formation energy of the unit cell and unit cell in-plane area, respectively. Using Eq. (4), we can calculate ΔGu for 5
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the ability to investigate the crystallographic orientations beyond the lowest-Miller-index surfaces, which could be analyzed in a following study. Nevertheless, the qualitative tendency is clear. With the change in the surface orientation from (100) to (111), the thin film growth rate decreases and the 2D growth approaches to a 3D island growth.
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4. Conclusion We investigated the crystallographic surface-orientation-dependent growth rate and growth modes of the SRO epitaxial thin films on the STO substrates using PLE. The growth characteristics were analyzed using XRD, AFM, and KPFM for the representative crystallographic surface orientations ((100), (110), and (111)) of the single-crystal thin films. With the change in the orientation from (100) to (110) to (111), the growth rate significantly decreased and the growth mode evolved from the 2D to the 3D growth. According to the first-principles calculation and atomistically modified nucleation barrier energy calculation, the unit-cell surface energy of SRO increased with the change in the orientation from (100) to (110) to (111), consistent with the experimentally observed changes. The systematic understanding of the singlecrystal thin films was further applied to the epitaxial polycrystalline thin film with arbitrarily oriented grains. Our study provides a design principle for functional oxide epitaxial thin films and heterostructures with different surface orientations necessary for conventional and future oxide opto-electronic applications. Acknowledgments This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF2019R1A2B5B02004546 and NRF-2019R1A6A1A03033215) and Programme Asie 127096 supported by the Conseil Régional de Normandie, and PHC Star 36453VL (AD & WP). Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.apsusc.2019.143924. References [1] G. Koster, L. Klein, W. Siemons, G. Rijnders, J.S. Dodge, C.-B. Eom, D.H.A. Blank, M.R. Beasley, Structure, physical properties, and applications of SrRuO3 thin films, Rev. Mod. Phys. 84 (2012) 253–298. [2] S.A. Lee, S. Oh, J.-Y. Hwang, M. Choi, C. Youn, J.W. Kim, S.H. Chang, S. Woo, J.S. Bae, S. Park, Y.-M. Kim, S. Lee, T. Choi, S.W. Kim, W.S. Choi, Enhanced electrocatalytic activity via phase transitions in strongly correlated SrRuO3 thin films, Energy Environ. Sci. 10 (2017) 924–930. [3] S.A. Lee, S. Oh, J. Lee, J.-Y. Hwang, J. Kim, S. Park, J.-S. Bae, T.E. Hong, S. Lee, S.W. Kim, W.N. Kang, W.S. Choi, Tuning electromagnetic properties of SrRuO3 epitaxial thin films via atomic control of cation vacancies, Sci. Rep. 7 (2017) 11583. [4] Y.J. Chang, C.H. Kim, S.H. Phark, Y.S. Kim, J. Yu, T.W. Noh, Fundamental thickness limit of itinerant ferromagnetic SrRuO3 thin films, Phys. Rev. Lett. 103 (2009) 057201. [5] D. Kan, Y. Shimakawa, Strain effect on structural transition in SrRuO3 epitaxial thin films, Cryst. Growth Des. 11 (2011) 5483–5487. [6] K. Ishigami, K. Yoshimatsu, D. Toyota, M. Takizawa, T. Yoshida, G. Shibata, T. Harano, Y. Takahashi, T. Kadono, V.K. Verma, V.R. Singh, Y. Takeda, T. Okane, Y. Saitoh, H. Yamagami, T. Koide, M. Oshima, H. Kumigashira, A. Fujimori, Thickness-dependent magnetic properties and strain-induced orbital magnetic moment in SrRuO3 thin films, Phys. Rev. B 92 (2015) 064402. [7] P. Kaur, K.K. Sharma, R. Pandit, R.J. Choudhary, R. Kumar, Structural, electrical, and magnetic properties of SrRuO3 thin films, Appl. Phys. Lett. 104 (2014) 081608.
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