Facile on-demand oriented growth of perovskite oxide thin films: applications of Dion–Jacobson phase as seed layer

Abstract

We investigated a new fabrication process for high-quality, uniaxial-oriented grown Dion–Jacobson perovskite (DJP) RbLaNb₂O₇ thin films by means of excimer laser assisted metal organic deposition. Oriented crystal growth occurred only by excimer laser irradiation in air at 400 °C after a deposition of metal–organic solution and preheating at 400 °C. The obtained RbLaNb₂O₇ thin film showed perfect (010)-oriented growth on borosilicate glass substrates, and the film surface had atomically flat terraces. We confirmed that the obtained DJP RbLaNb₂O₇ thin films functioned very well as a seed layer for the fabrication of various perovskite oxide thin films with high orientation quality. An oriented grown LaNiO₃ thin film fabricated on the DJP seed layer and prepared at 400 °C exhibited low electrical resistivity at room temperature. Moreover, we demonstrated the utility of our methodology for on-demand fabrication by combining our method with an inkjet printing technique to pattern LaNiO₃ thin films with a line shape, which also were successfully grown with uniaxial orientation on the DJP seed layer. A ferroelectric Na_0.5K_0.5NbO₃ perovskite film was heteroepitaxially grown on the uniaxial-oriented grains of the LaNiO₃/RbLaNb₂O₇ thin film, and this oriented film exhibited a higher dielectric constant than that of a film fabricated without the DJP seed layer. Thus, we demonstrated facile, oriented crystal growth enabled by the uniaxial-oriented DJP seed layer.

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Introduction

Physical properties of oxide thin films are strongly dependent on their crystal quality. In particular, control of crystal orientation is considered to be one of the most important issues in obtaining good properties of oxide thin film materials, and to date such control has generally been achieved by designing surface structures on substrates to promote specific crystal orientations. For many studies of oxide thin films, single crystal oxide substrates such as SrTiO₃, LaAlO₃, α-Al₂O₃, and MgO have been used because of their precise surface structural control that enables high-quality crystal growth in particular orientations.^1–5 However, the use of single crystal oxide substrates is not realistic for industrial applications due to the crystals' high cost and difficulty in up scaling. Therefore, crystal growth processes for high-quality oriented oxide thin films without single crystal substrates have been recently developed; some seed layer methods, such as ion beam assisted deposition and Langmuir–Blodgett nanosheet deposition.^6–10 Though these processes have enabled the fabrication of high-quality, oriented oxide heterostructures, further improvements are still needed to allow on-demand fabrication of oriented films for oxide electronics applications, such as microelectromechanical systems (MEMS),^11,12 owing to process restrictions associated with these applications. Thus, further development of oriented film growth techniques is an important issue for future high-integration applications.

In this study, we report a new process based on chemical solution deposition (CSD) for preparation of perovskite oxide multi-layered thin films with high-quality uniaxial orientation. The CSD process was very simple and enabled on-demand fabrication, in contrast to the conventional methods. To achieve the objectives, we employed an excimer laser-assisted metal organic deposition (ELAMOD) process developed from a conventional metal organic deposition (MOD) process.^13–16 The ELAMOD process can be used to fabricate oxide thin films at low temperature for commonly used substrates such as Si and borosilicate glass, which cannot withstand high temperatures (T > 400 °C) in air: oxide films are crystallized by means of excimer laser irradiation (λ = 193, 248, and 308 nm) instead of by high-temperature furnace heating as is used in the MOD process. Details of the ELAMOD crystal growth mechanism are described elsewhere.^15,16 In this study, we chose the Dion–Jacobson perovskite (DJP) oxide RbLaNb₂O₇^17–19 to prepare an oriented growth seed layer on amorphous glass substrates by means of the ELAMOD process. The crystal structure of RbLaNb₂O₇ (referred to herein as “DJP seed layer”) is shown in Fig. 1.¹⁹ The Rb ions and double-layered perovskite slabs are stacked alternately along the b-axis. We investigated the oriented growth of DJP seed layer thin films on amorphous glass substrates, and we demonstrate herein the fabrication of various multi-layered perovskite thin films on the obtained oriented RbLaNb₂O₇ seed layers.

Fig. 1 Schematic crystal structure of RbLaNb₂O₇ used to fabricate DJP seed layer films. Rb ions and double-layered perovskite slabs are stacked alternately along the b-axis.

Experimental

Thin film fabrication procedures

DJP seed layer (RbLaNb₂O₇) thin films were fabricated by the ELAMOD process. The starting solutions for the RbLaNb₂O₇ films were prepared by mixing 2-ethylhexanoate solutions of the constituent metals diluted with toluene to obtain the required concentration and viscosity for spin-coating. The Rb : La : Nb molar ratio in the coating solution was 1.0 : 1.0 : 2.0. This solution was spin-coated onto borosilicate glass substrates at 4000 rpm for 10 s. The coated films were preheated at 400 °C in air for 10 min to decompose the organic components of the films. The spin coating and preheating were repeated three times, and then the films were irradiated with a KrF laser (Lambda Physik Compex110) at an afluence of 60 mJ cm⁻² and a pulse duration of 26 ns for 7500 pulses at 400 °C in air. The laser energy was homogenized by a beam homogenizer. Thereafter, the obtained RbLaNb₂O₇ thin films were used as a DJP seed layer for the preparation of various perovskite materials as described below. SrTiO₃ thin films on borosilicate glass substrates with and without the DJP seed layer were prepared by means of the ELAMOD process. A mixed solution of strontium and titanium 2-ethylhexanoates (Sr : Ti = 1.0 : 1.0) diluted with toluene was spin-coated onto substrates at 4000 rpm for 10 s. The coated films were dried at 100 °C in air to remove the solvent, and then preheated at 400 °C in air for 10 min. The preheated films were irradiated with the KrF laser at an afluence of 75 mJ cm⁻² for 7500 pulses at 400 °C in air. For conducting LaNiO₃ thin films on borosilicate glass substrates with and without the DJP seed layer, lanthanum and nickel 2-ethylhexanoate solutions (La : Ni = 1.0 : 1.0) diluted with toluene were used as a starting solution. The solution was spin-coated onto substrates at 4000 rpm for 10 s. In addition, a line-shaped LaNiO₃ thin film was also prepared on the DJP seed layer by deposition of the lanthanum–nickel starting solution using an inkjet printer (Microjet, FemtoJet). The deposited films were dried at 100 °C in air and then preheated at 400 °C in air for 10 min. The preheated films were irradiated with the KrF laser at an afluence of 50 mJ cm⁻² for 5000 pulses at 400 °C in air. On the obtained bottom electrode LaNiO₃ thin films, ferroelectric Na_0.5K_0.5NbO₃ thin films were fabricated. The Na : K : Nb molar ratio in the coating solution was 1.0 : 1.0 : 2.0. This solution was spin-coated onto the LaNiO₃ thin films at 4000 rpm for 10 s. The coated films were dried at 100 °C in air and preheated at 400 °C in air for 10 min. The preheated films were irradiated with the KrF laser at an afluence of 50 mJ cm⁻² for 5000 pulses at 400 °C in air. The film thickness of each perovskite thin film on the seed layer was increased by repeating the coating, preheating, and laser irradiation processes. All of the processes were carried out at low temperature (<400 °C) in air.

Sample characterization

The crystal structure and orientation properties of the obtained films were studied by X-ray diffraction measurements (Rigaku, SmartLab). 2θ − β maps (β: the direction along Debye rings), a χ − ϕ map, and micron-scale analysis were obtained with a two-dimensional (2D) pixel area detector (Dectris, PILATUS 100 K)²⁰ and a collimator with 200 μm diameter. The detector length from the samples was fixed at 102 mm. The degree of crystal orientation was evaluated in terms of the Lotgering factor F, which is calculated from the following equation.²¹

F(hkl) = (P − P₀)/(1 − P₀)

where P₀ = ΣI₀(hkl)/ΣI₀(HKL) and P = ΣI(hkl)/ΣI(HKL). I₀ and I are the intensities of each of the diffraction peaks in X-ray diffraction patterns as presented in ICSD database and in experimental data, respectively. The relative growth rate of the crystal planes was determined for each facet area of the RbLaNb₂O₇ as calculated by the Bravais–Friedel Donnay–Harker (BFDH) method²² using a Materials Studio program package. The crystal growth process and crystallinity of the thin films were observed by cross-sectional transmission electron microscopy (XTEM) and selected area electron diffraction (SAED) analysis using a JEM-2010 (JEOL) instrument operating at 200 kV. The samples for the XTEM observations were prepared by argon milling using an ion milling device (BAL-TEC, RES100). The surface of the samples was studied with an atomic force microscope (AFM). The electrical resistivity curves of the LaNiO₃ thin films as a function of temperature were measured by a four-probe technique. The dielectric constants of the Na_0.5K_0.5NbO₃ thin films were collected at 1 kHz with a LCR meter (WAYNE KERR, 6440A).

Results and discussion

Oriented growth of Dion–Jacobson perovskite RbLaNb₂O₇ thin films

Fig. 2 shows the 2θ/ω scan of X-ray diffraction for a ELAMOD-fabricated DJP RbLaNb₂O₇ film with a thickness of 60 nm on a glass substrate. The RbLaNb₂O₇ film was very well crystallized without impurities, and all of the observed Bragg reflections were assigned to 0k0 indices without any other orientations. This assignment indicates that the film was grown with a strong (010) orientation, corresponding to the layered stacking direction of the DJP structure shown in the inset of Fig. 2. From the X-ray diffraction intensities, the Lotgering factor F(010) was estimated to be 1.0, indicating a perfect (010) orientation. To analyze the orientation quality of the obtained film, we evaluated 2D X-ray diffraction measurements. The 2θ − β map for the RbLaNb₂O₇ film is shown in Fig. 3a. The Bragg reflections corresponding to 0k0 indices appeared as spots at β = 0°. The β scan profiles for 020 and 080 reflections were collected from the obtained 2θ − β map. Both reflection profiles showed an obvious peak shape, and their full widths at half maximum (FWHM) were 3.70° and 3.85°, respectively. These values can be regarded as small for the crystal orientation quality on the amorphous glass substrates. Fig. 3c shows the χ − ϕ map around the 141 reflection. The 141 reflection was observed along χ = 55° at all values of ϕ. The obtained results indicate that the in-plane crystallized grains of the film had random orientations, although the in-plane crystal grain size appeared to be relatively large as evidenced by the inhomogeneous distribution of peak intensities. On the basis of these X-ray diffraction measurements, we confirmed that the fabricated RbLaNb₂O₇ film exhibited high-quality, uniaxial, (010)-oriented crystal growth.

Fig. 2 X-Ray diffraction pattern for a DJP RbLaNb₂O₇ thin film on a borosilicate glass substrate (top), and the simulated pattern of polycrystalline RbLaNb₂O₇ with random orientation (bottom). The inset shows a schematic illustration for the relationship between the obtained film orientation and the substrate.

Fig. 3 (a) 2θ − β map for a DJP RbLaNb₂O₇ thin film on a borosilicate glass substrate. (b) β scan profiles collected from the 2θ − β map at 2θ = 8.02° and 32.72° for 020 and 080 reflections, respectively. (c) χ − ϕ map around the 141 reflection.

Fig. 4 shows a high-resolution XTEM image (a) and the SAED pattern (b) of the prepared RbLaNb₂O₇ film along the [101] projection. The layered structure corresponding to the alternate stacking of Rb ions and double-layered perovskite slabs along the b-axis of RbLaNb₂O₇ was clearly observed in a direction normal to the substrate surface. The emerging stacking faults observed in the structure shown in Fig. 4a probably gave rise to the larger FWHM (ca. 3.70°) observed in β-scans of 0k0 X-ray diffraction peaks of the RbLaNb₂O₇ film compared to those of peaks generally observed in single crystals. Fig. 5a shows the XTEM image of the sample after the excimer laser irradiation at 3000 pulses to the amorphous precursor film. A crystalline, layered phase about 20 nm was confirmed at the film surface, and a partially crystalline phase was observed under the crystalline phase. The location of these phases indicates that the crystal growth of the RbLaNb₂O₇ thin film preferentially proceeded from the amorphous film surface to the interface with the glass substrate.

Fig. 4 (a) [101]-zone high-resolution XTEM image for a DJP RbLaNb₂O₇ thin film on a borosilicate glass substrate after excimer laser irradiation at 7500 pulses. The white arrows represent stacking faults of the DJP layered structure. (b) SAED pattern of the RbLaNb₂O₇ thin film.

Fig. 5 (a) [101]-zone high-resolution XTEM image for a DJP RbLaNb₂O₇ thin film after excimer laser irradiation at 3000 pulses, and brightness analysis of the obtained XTEM image as a function of depth. The brightness periodicity indicates the presence of a crystalline DJP phase, and one period of brightness intensity corresponds to one layer unit (1 l.u.) of the DJP structure as shown in the inset. (b) Schematic illustrations of the depth dependence of the temperature distribution under pulsed laser irradiation, and of the crystal growth process with increasing number of laser pulses.

Two remarkable features associated with crystal growth of oxide films in the ELAMOD process should be noted here. First, when no effective crystal nucleation sites are present in an amorphous precursor film matrix, crystal nucleation occurs from the film surface, since the excimer laser irradiation induces a gradient temperature profile through a film thickness of several hundred nanometres due to a large photoabsorbance of the films (left panel of Fig. 5b).¹⁶ Second, epitaxial growth proceeds preferentially from active nucleation sites, provided that the sites have a large absorbance for the irradiated laser wavelength and a small lattice mismatch with the grown material.¹⁵ Therefore, the initial nucleation of crystalline RbLaNb₂O₇ on the glass substrate probably occurred at the amorphous precursor film surface under excimer laser irradiation. The film then would have grown epitaxially from those first nucleation sites at the film surface, as the crystallized RbLaNb₂O₇ has a large absorbance at the KrF excimer laser wavelength (248 nm). In addition, the potential feature of the crystal growth in RbLaNb₂O₇ is essential for the orientation growth in this process. Crystal facet area calculations associated with the crystal growth rates (Table 1) revealed that the {020} plane was largely preferred among the crystal facets of RbLaNb₂O₇, reflecting the two-dimensionality of the DJP layered structure. This suggests that the (010) orientation at the film surface would significantly reduce the surface energy in the crystal growth process. Hence, the growth mechanism of (010) orientation corresponding to the DJP layered structure in this process can be explained as follows. First, nucleation owing to the gradient temperature distribution induced by the excimer laser irradiation occurred at the film surface, while the (010)-layered structural ordering reduced the crystal surface energy. Then, the film grew epitaxially from the first (010)-oriented crystal nuclei toward the interface with the borosilicate glass substrate. Thus, we concluded that the uniaxial-oriented growth observed in this study was enabled by the complementary nature of the ELAMOD process to the DJP RbLaNb₂O₇ crystal growth.

Table 1 Relative stability of crystal facets for the DJP RbLaNb₂O₇

{hkl}	d(hkl)/Å	Total facet area (%)
020	10.995	54.03
011	5.329	19.38
101	3.884	26.58
200	2.747	0.01

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The surface of the prepared RbLaNb₂O₇ film was observed by AFM (Fig. 6a), and the root-mean-squared surface roughness of the film was evaluated to be 9.46 nm. This value is large compared to that of commercial single crystal substrates. We conducted additional evaluations of the surface roughness of the film. Fig. 6b shows the XTEM image of the film around the variable thickness region that caused the surface roughness observed in the AFM image. It should be noted that the film surface in the flat region of the film (labelled 1 in Fig. 6b and 6c) was confirmed to be atomically flat, comparable to the surfaces of perovskite nanosheets and cleaved layered crystals.^9,23–26 This flatness would have originated from the reduced crystal surface energy associated with the crystal growth of RbLaNb₂O₇ discussed above. Moreover, remarkable features were also observed in regions of varied thickness in the film (labelled 2 and 3 in Fig. 6b and 6c), including steps the size of layer units (the thickness of one layer unit is about 1.1 nm) and atomically flat terraces without ambiguous disordered lattice morphology. These results mean that the RbLaNb₂O₇ film had a clean (010) crystal surface, which is favorable for seed layers used to grow uniaxial-oriented perovskite heterostructures, at almost surface area, even though the film had a relatively large surface roughness. Thus, these results demonstrated perfect (010)-oriented growth of the DJP RbLaNb₂O₇ thin film at low temperature by means of the CSD-based ELAMOD process that we developed. In particular, we note that this process was extremely facile, involving only three steps: coating, preheating, and laser irradiation.

Fig. 6 (a) AFM image of the surface of a DJP RbLaNb₂O₇ thin film. (b) XTEM image of the thin film after excimer laser irradiation at 7500 pulses, and (c) high-resolution XTEM images of the enlarged views for the regions labelled 1, 2, and 3 in (b). The black arrows indicate atomically flat surface regions. The red indicators represent one layer unit of the DJP structure.

Seed layer property of RbLaNb₂O₇ thin films for uniaxial-oriented growth of perovskites

We studied the utility of the obtained (010)-oriented DJP RbLaNb₂O₇ thin film as a seed layer for preparation of uniaxial-oriented perovskite oxide thin films. We prepared thin films of SrTiO₃, which is a representative cubic perovskite oxide with lattice parameter a = 3.905 Å (the lattice mismatch of SrTiO₃ with the b-plane of RbLaNb₂O₇ in the simple perovskite units is 0.53%),²⁷ on glass substrates with and without the DJP seed layer by means of the ELAMOD process at 400 °C in air. The 2θ/ω scans of X-ray diffraction of the prepared SrTiO₃ thin films with 60 nm thickness are shown in Fig. 7. On the bare glass substrate, the observed diffraction peaks were indexed to 100, 110, and 200 for the cubic SrTiO₃, indicating almost random orientations: the Lotgering factor F(100) was evaluated to be 0.268. In contrast, the SrTiO₃ film on the seed layer was well crystallized and showed a highly (100)-oriented growth. The main diffraction peaks were assigned to h00, whereas a trace of 110 peak intensity was also confirmed near 32.5° by multi-peak fittings. The Lotgering factor F(100) was calculated to be 0.983, which was significantly large.

Fig. 7 X-Ray diffraction patterns for prepared films of RbLaNb₂O₇/glass (blue), SrTiO₃/RbLaNb₂O₇/glass (red), and (c) SrTiO₃/glass (black). The inset shows an enlarged view of each diffraction pattern at 31.5–34.0°. The subscripts R and S represent RbLaNb₂O₇ and SrTiO₃, respectively.

Fig. 8a and 8b show the 2θ − β mappings for the obtained SrTiO₃ films without and with the seed layer, respectively. Although the Bragg reflections of the SrTiO₃ film without the seed layer exhibited spread Debye rings along β, those for the SrTiO₃ film with the seed layer were clearly observed as defined spots. The β-scans for the 080 diffraction of the seed layer and 200 diffraction of the SrTiO₃ film are shown in Fig. 8c. The FWHM of these peaks were 3.72° and 5.52°, respectively. These values suggest that the SrTiO₃ film was well crystallized and oriented in the (100) direction with a nearly equivalent orientation quality in the seed layer. The slight deterioration of the FWHM value indicates that the increasing orientation distribution might have been caused by the imperfect surface flatness of the seed layer.

Fig. 8 2θ − β map for (a) SrTiO₃/glass and (b) SrTiO₃/RbLaNb₂O₇/glass. (c) β-scan profiles collected from the 2θ − β map for 080_R (2θ = 32.72°) and 200_S (2θ = 46.48°) reflections in SrTiO₃/RbLaNb₂O₇/glass. The subscripts R and S represent RbLaNb₂O₇ and SrTiO₃, respectively.

Fig. 9 shows the XTEM and SAED analyses for the prepared (100)-oriented SrTiO₃ film grown on the DJP seed layer. The high-resolution lattice images and SAED patterns of both phases showed that the SrTiO₃ film was grown epitaxially on the (010)-oriented grain of the seed layer: the cube-on-cube relationship was clearly maintained at the interface between the flat surface of the seed layer and the SrTiO₃ film, suggesting that the (010)-oriented RbLaNb₂O₇ functioned well as the seed layer for the fabrication of perovskite oxide, uniaxial-oriented heterostructures.

Fig. 9 (a) High resolution XTEM image of a prepared multi-layered SrTiO₃/RbLaNb₂O₇ thin film on a borosilicate glass substrate. The enlarged lattice images and SAED patterns for (b) the SrTiO₃ film and (c) the RbLaNb₂O₇ film are also shown. The subscript P represents the unit cell for the simple cubic perovskite.

Application of the DJP seed layer for conducting and ferroelectric perovskite materials

The pseudo-cubic perovskite LaNiO₃ with a = 3.840 Ų⁸ (the lattice mismatch is 1.15% with RbLaNb₂O₇ in the simple perovskite units) is well known as a material for conducting electrodes.^29–33 Polycrystalline LaNiO₃ thin films have exhibited electrical resistivity characteristic of a metallic material, about 0.5–2.0 mΩ cm, at room temperature.^30–33 However, a relatively high process temperature (about 700–800 °C) was used to produce these metallic films. Therefore, we prepared metallic LaNiO₃ films by means of the ELAMOD process with the DJP RbLaNb₂O₇ seed layer at a low process temperature (below 400 °C).

Fig. 10 shows the 2θ/ω scans of X-ray diffraction for the LaNiO₃ films prepared at 400 °C in air with a thickness of ca. 100 nm with and without the DJP seed layer. The crystallized LaNiO₃ film on the bare glass substrate showed random orientation growth without (100) orientation. In contrast, the LaNiO₃ film on the seed layer was well crystallized and exhibited a significant enhancement of h00 diffraction peaks, which indicated the (100) orientation. The Lotgering factor F(100) of this film was calculated to be 0.971. The electrical resistivities (ρ) of the prepared LaNiO₃ films with and without the DJP seed layer at room temperature as a function of film thickness are shown in Fig. 11. The ρ values for both samples decreased with increasing film thickness up to about 60 nm. At a film thickness of about 100 nm, ρ of the obtained LaNiO₃ films with and without the DJP seed layer were 4.42 and 1.02 mΩ cm, respectively. Hence, the film with the DJP seed layer exhibited a 77% decrease in resistivity compared to the film on the bare substrate, and the temperature dependence of ρ was good metal to low temperature. Furthermore, the ρ value of the LaNiO₃ film on the DJP seed layer was comparably lowered to that of the films prepared at high process temperatures in the previous reports.^29–33

Fig. 10 X-Ray diffraction patterns for prepared films of LaNiO₃/RbLaNb₂O₇/glass (red) and LaNiO₃/glass (black). The inset shows the enlarged view of each diffraction pattern at 31.5–34.0°. The subscripts R and L represent RbLaNb₂O₇ and LaNiO₃, respectively.

Fig. 11 ρ for LaNiO₃/RbLaNb₂O₇/glass and LaNiO₃/glass as a function of LaNiO₃ film thickness. The inset shows the temperature dependence of ρ for LaNiO₃/RbLaNb₂O₇/glass with a LaNiO₃ thickness of ca. 100 nm.

Fig. 12a displays an optical microscope image of the line-shape patterned LaNiO₃ film on the DJP seed layer after ELAMOD processing. The starting solution was deposited by an inkjet printer (see Experimental). The width and thickness of the patterned LaNiO₃ wire film were evaluated to be 60 μm and 30 nm, respectively (inset of Fig. 12). Fig. 12b shows the 2θ-ω scan of the region including the patterned LaNiO₃ film by using a ϕ200 μm X-ray collimator. All of the observed Bragg reflections were assigned to h00 peaks for LaNiO₃ and 0k0 peaks for the DJP seed layer, indicating the successful (100)-oriented growth of the microscopic patterned LaNiO₃ film. The compatibility of the ELAMOD process with the inkjet printing technique and with the DJP seed layer indicates that our developed method could be used for on-demand fabrication of uniaxial oriented perovskite oxide thin films, which would be very useful for integrated devices such as MEMS.

Fig. 12 (a) Optical microscope image for a patterned LaNiO₃ film on the DJP seed layer after excimer laser irradiation. The inset shows an AFM image for the patterned LaNiO₃ film. (b) Microscopic analysis of X-ray diffraction for the patterned LaNiO₃ film on the DJP seed layer. The inset shows the β-scan profile for the 200_L (2θ = 47.46°) reflection. The subscripts R and L represent RbLaNb₂O₇ and LaNiO₃, respectively.

We also observed the heteroepitaxial multi-layered structure of perovskite films on uniaxial oriented grains of the DJP seed layer. Ferroelectric Na_0.5K_0.5NbO₃ thin films with a thickness of ca. 420 nm were fabricated on the aforementioned LaNiO₃ thin films as a bottom electrode with and without the DJP seed layer by means of the ELAMOD process. Fig. 13 shows the 2θ/ω scans of X-ray diffraction for the Na_0.5K_0.5NbO₃ thin films on the LaNiO₃ electrode films with and without the seed layer. On the LaNiO₃ thin film with random orientations, the Na_0.5K_0.5NbO₃ thin film exhibited a weak (100) orientation; the Lotgering factor F(100) was calculated to be 0.552. In contrast, the Na_0.5K_0.5NbO₃ thin film fabricated on the highly (100)-oriented LaNiO₃ thin film was crystallized with uniaxial orientation: the Lotgering factor F(100) in this case was calculated to be 0.943. This result indicates the utility of the DJP seed layers for fabricating multi-layered, uniaxial-oriented growth of perovskite heterostructures. The small decrease in the uniaxial orientation of the Na_0.5K_0.5NbO₃ thin film compared to that of the (100)-oriented LaNiO₃ thin film might have been caused by the relatively large lattice mismatch (about 2.79%) between the films' respective perovskite unit cells.³⁴

Fig. 13 X-Ray diffraction patterns for prepared films of Na_0.5K_0.5NbO₃/LaNiO₃/RbLaNb₂O₇/glass (red) and Na_0.5K_0.5NbO₃/LaNiO₃/glass (black). The inset shows the temperature dependence of the relative dielectric constants for Na_0.5K_0.5NbO₃ films with (red) and without (black) the DJP seed layer. The subscripts R, L, and N represent RbLaNb₂O₇, LaNiO₃, and Na_0.5K_0.5NbO₃, respectively.

The relative dielectric constants (ε_r) of the prepared samples at room temperature were measured by applying AC 0.1 V at 1 kHz frequency. The ε_r values for the Na_0.5K_0.5NbO₃ thin films on the (100)-oriented and non-oriented LaNiO₃ electrodes were 570 and 325, respectively. As shown in the inset of Fig. 13, enhancement of ε_r in the highly crystallized and oriented grown film was clearly observed over a wide temperature range including the two structural phase transition temperatures of Na_0.5K_0.5NbO₃ near 218 °C (orthorhombic–tetragonal) and 420 °C (tetragonal–cubic)^35–37 owing to the high crystal orientation degree and the grain growth effect as observed in the previous reports.^38–40 The broader temperature dependence of ε_r in the obtained Na_0.5K_0.5NbO₃ films compared to that of the ceramic samples could be caused by multiple phenomena, such as the strain effect from the substrate lattices, the thermal excitation of dipolar (double ionized oxygen vacancies and niobium ions) under the alternative electric field,⁴¹ and the smaller crystal grain size of the films compared to that of the ceramic samples.

Therefore, we demonstrated that the DJP seed layer prepared by the ELAMOD process was successfully utilized for the fabrication of uniaxial-oriented, multi-layered perovskite heterostructures. All of the fabrication process temperatures for the various perovskite oxide thin films were low, below 400 °C, and we also demonstrated that this process was suitable in combination with an inkjet printing technique for the on-demand fabrication of uniaxial-oriented grown perovskite oxide thin films.

Conclusions

A new fabrication process for high-quality, uniaxial-oriented grown DJP RbLaNb₂O₇ thin films by means of an ELAMOD process has been investigated: oriented crystal growth occurred only by excimer laser irradiation in air at 400 °C after deposition of a metal–organic solution and preheating at 400 °C. The obtained RbLaNb₂O₇ thin film exhibited perfect (010)-oriented growth on borosilicate glass substrates, and the film surface had atomically flat terraces. We confirmed that the characteristics of the obtained DJP thin films functioned very well as a seed layer for the fabrication of various perovskite oxide thin films with high orientation quality. An oriented grown LaNiO₃ thin film prepared on the DJP seed layer at 400 °C exhibited a low electrical resistivity of 1.02 mΩ cm at room temperature. Moreover, we demonstrated the utility of our methodology for on-demand fabrication: a patterned LaNiO₃ thin film deposited as a line shape by inkjet printing was successfully grown with uniaxial orientation on the DJP seed layer. A ferroelectric Na_0.5K_0.5NbO₃ perovskite film was also heteroepitaxially grown on the uniaxial-oriented grains of the LaNiO₃/RbLaNb₂O₇ thin film, and the oriented film showed a higher dielectric constant than that of a film grown without the seed layer. Thus, we demonstrated facile, oriented crystal growth enabled by the uniaxial-oriented DJP seed layer.

Acknowledgements

We are grateful to M. Ichihara at the Institute for Solid State Physics, the University of Tokyo, for his help as well as valuable discussions. The XTEM observations and electrical resistivity measurements were carried out at the facilities of the Materials Design and Characterization Laboratory, Institute for Solid State Physics, University of Tokyo.

References

1. A. Brinkman, M. Huijben, M. Vanzalk, J. Huijben, U. Zeitler, J. C. Maan, W. G. van der Wiel, G. Rijnders, D. H. A. Blank and H. Hilgenkamp, Nat. Mater., 2007, 6, 493.
2. I. Vrejoiu, M. Alexe, D. Hesse and U. Gösele, Adv. Funct. Mater., 2008, 18, 3892.
3. S. J. May, P. J. Ryan, J. L. Robertson, J.-W. Kim, T. S. Santos, E. Karapetrova, J. L. Zarestky, X. Zhai, S. G. E. te Velthuis, J. N. Eckstein, S. D. Bader and A. Bhattacharya, Nat. Mater., 2009, 8, 892.
4. Y. Wakabayashi, D. Bizen, H. Nakao, Y. Murakami, M. Nakamura, Y. Ogimoto, K. Miyano and H. Sawa, Phys. Rev. Lett., 2006, 96, 017202.
5. S. A. Chambers, Adv. Mater., 2009, 21, 1622.
6. M. Q. Huang, J. Greek, S. Massing, O. Meyer, H. Reiner and G. Linker, Nucl. Instrum. Methods Phys. Res., Sect. B, 1999, 148, 793.
7. T. Kato, Y. Iijima, T. Muroga, T. Saitoh, T. Hirayama, I. Hirabayashi, Y. Yamada, T. Izumi, Y. Shiohara and Y. Ikuhara, Phys. C, 2003, 392–396, 790.
8. K. Kikuta, K. Noda, S. Okumura, S. Yamaguchi and S. Hirano, J. Sol-Gel Sci. Technol., 2007, 42, 381.
9. T. Shibata, K. Fukuda, Y. Ebina, T. Kogure and T. Sasaki, Adv. Mater., 2008, 20, 231.
10. T. Shibata, T. Ohnishi, I. Sakaguchi, M. Osada, K. Takada, T. Kogure and T. Sasaki, J. Phys. Chem. C, 2009, 113, 19096.
11. M. Heule, S. Vuillemin and L. J. Gauckler, Adv. Mater., 2003, 15, 1237.
12. S. W. Kirchoefer, E. J. Cukauskas, N. S. Barker, H. S. Newman and W. Chang, Appl. Phys. Lett., 2002, 80, 1255.
13. T. Tsuchiya, K. Daoudi, T. Manabe, I. Yamaguchi and T. Kumagai, Appl. Surf. Sci., 2007, 253, 6504.
14. T. Tsuchiya, I. Yamaguchi, T. Manabe, T. Kumagai and S. Mizuta, Mater. Sci. Semicond. Process., 2002, 5, 207.
15. T. Nakajima, T. Tsuchiya, M. Ichihara, H. Nagai and T. Kumagai, Chem. Mater., 2008, 20, 7344.
16. T. Nakajima, T. Tsuchiya, M. Ichihara, H. Nagai and T. Kumagai, Appl. Phys. Express, 2009, 2, 023001.
17. M. Dion, M. Ganne and M. Tournoux, Mater. Res. Bull., 1981, 16, 1429.
18. J. Gopalakrishnan, V. Bhat and B. Raveau, Mater. Res. Bull., 1987, 22, 413.
19. A. R. Armstrong and P. A. Anderson, Inorg. Chem., 1994, 33, 4366.
20. P. Kraft, A. Bergamaschi, Ch. Broennimann, R. Dinapoli, E. F. Eikenberry, B. Henrich, I. Johnson, A. Mozzanica, C. M. Schlepütz, P. R. Willmott and B. Schmitt, J. Synchrotron Radiat., 2009, 16, 368.
21. F. K. Lotgering, J. Inorg. Nucl. Chem., 1959, 9, 113.
22. J. D. H. Donnay and D. Harker, Amer. Mineralogist, 1937, 22, 446.
23. M. M. Treacy, S. B. Rice, A. J. Jacobson and J. T. Lewandowski, Chem. Mater., 1990, 2, 279.
24. R. E. Schaak and T. E. Mallouk, Chem. Mater., 2000, 12, 3427.
25. Y. Wakabayashi, M. H. Upton, S. Grenier, J. P. Hill, C. S. Nelson, J.-W. Kim, P. J. Ryan, A. I. Goldman, H. Zheng and J. F. Mitchell, Nat. Mater., 2007, 6, 972.
26. R. Matzdorf, Ismail, T. Kimura, Y. Tokura and E. W. Plummer, Phys. Rev. B: Condens. Matter Mater. Phys., 2002, 65, 085404.
27. R. J. Nelmes, G. M. Meyer and J. Hutton, Ferroelectrics, 1978, 21, 461.
28. A. Wold, B. Post and E. Banks, J. Am. Chem. Soc., 1957, 79, 4911.
29. K. Sreedhar, J. Honig, M. Darwin, M. McElfresh, P. Shand, J. Xu, B. Crooker and J. Spalek, Phys. Rev. B: Condens. Matter, 1992, 46, 6382.
30. A.-D. Li, C.-Z. Ge, D. Wu, P. Lü, Y.-Q. Zuo, S.-Z. Yang and N.-B. Ming, Thin Solid Films, 1997, 298, 165.
31. K. Ueno, W. Sakamoto, T. Yogo and S. Hirano, Jpn. J. Appl. Phys., 2001, 40, 6049.
32. D. Bao, X. Yao, N. Wakiya, K. Shinozaki and N. Mizutani, J. Phys. D: Appl. Phys., 2003, 36, 1217.
33. L. Yang, G. Wang, C. Mao, Y. Zhang, R. Liang, C. Soyer, D. Rémiens and X. Dong, J. Cryst. Growth, 2009, 311, 4241.
34. T. Saito, T. Wada, H. Adachi and I. Kanno, Jpn. J. Appl. Phys., 2004, 43, 6627.
35. Y. P. Guo, K. Kakimoto and H. Ohsato, Appl. Phys. Lett., 2004, 85, 4121.
36. B. Zhang, L. Zhang, J. Li, X. Ding and H. Zhang, Ferroelectrics, 2007, 358, 188.
37. D. Lin, K. W. Kwok and H. L. W. Chan, Appl. Phys. Lett., 2007, 90, 232903.
38. C.-R. Cho and A. Grishin, Appl. Phys. Lett., 1999, 75, 268.
39. H. Takao, Y. Saito, Y. Aoki and K. Horibuchi, J. Am. Ceram. Soc., 2006, 89, 1951.
40. K. Shibata, K. Suenaga, A. Nomoto and T. Mishima, Jpn. J. Appl. Phys., 2009, 48, 121408.
41. J. Miao, X. G. Xu, Y. Jiang, L. X. Cao and B. R. Zhao, Appl. Phys. Lett., 2009, 95, 132905.

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