Control of Polar Orientation and Lattice Strain in Epitaxial BaTiO

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Functional Inorganic Materials and Devices

Control of Polar Orientation and Lattice Strain in Epitaxial BaTiO Films on Silicon 3

Jike Lyu, Saúl Estandía, Jaume Gazquez, Matthew F. Chisholm, Ignasi Fina, Nico Dix, Josep Fontcuberta, and Florencio Sanchez ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b07778 • Publication Date (Web): 09 Jul 2018 Downloaded from http://pubs.acs.org on July 10, 2018

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ACS Applied Materials & Interfaces

Control of Polar Orientation and Lattice Strain in Epitaxial BaTiO3 Films on Silicon Jike Lyu,1 Saúl Estandía,1 Jaume Gazquez,1 Matthew F. Chisholm,2 Ignasi Fina,1 Nico Dix,1 Josep Fontcuberta,1 Florencio Sánchez*,1 1

Institut de Ciència de Materials de Barcelona (ICMAB-CSIC), Campus UAB, Bellaterra 08193, Barcelona, Spain 2

Materials Science and Technology Division, Oak Ridge National Laboratory, Tennessee 37831-6071, USA

*Email: [email protected] KEYWORDS: ferroelectric oxides; epitaxial oxides on silicon; strain engineering; barium titanate; pulsed laser deposition; dipoles maps ABSTRACT: Conventional strain engineering of epitaxial ferroelectric oxide thin films is based on the selection of substrates with suitable lattice parameter. Here we show that the variation of oxygen pressure during pulsed laser deposition is a flexible strain engineering method for epitaxial ferroelectric BaTiO3 films either on perovskite substrates or on Si(001) wafers. This unconventional growth strategy permits continuous tuning of strain up to high levels (ε>0.8%) in films of thickness exceeding of a hundred of nanometers, as well as selecting the polar axis orientation to be either parallel or perpendicular to the substrate surface plane.

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1. INTRODUCTION The properties of functional oxides can be modified by small variations in their crystal structure. In case of epitaxial films the lattice mismatch induces structural distortions, permitting the control of the functional properties when a specific substrate is chosen. This strain engineering has been widely applied to superconducting, metallic, ferromagnetic, and ferroelectric oxides, among others.1-6 For example, the ferroelectric polarization and Curie temperature of ferroelectric BaTiO3 (BTO) epitaxial films compressively strained on scandate substrates can be substantially enhanced.7 Conventional strain engineering is based on selecting a substrate with a particular lattice mismatch with the film, and this allows a range of discrete values for lattice deformation restricted to the critical thickness below the onset of plastic relaxation. Moreover, many current applications require integration with silicon wafers, thus losing the possibility of selecting a specific substrate. Thus, conventional strain engineering is limited to either films very thin, or deposited on substrates presenting low lattice mismatch (thus causing low strain). This last restriction is crucial if, for any reason, the substrate must be Silicon. Therefore alternative strategies to the classic strain engineering, with flexibility to be used on different substrates as silicon and not restricted to ultrathin films, have to be developed. Defects in thin films8-12 if controlled properly, offer an opportunity to engineer the strain using growth kinetics to expand the crystal lattice and control defects concentration. The crystal lattice of films deposited on a particular substrate, for example silicon, can be tuned using deposition parameters. For instance, growth rate and deposition temperature in pulsed laser deposition (PLD) of SrTiO3 (STO) films13 and ferroelectric BTO/STO superlattices14 and BTO films15 have been shown to permit control lattice strain.


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For ferroelectric oxides, the control of polar axis orientation (in-plane or out-of-plane) is also critical. The crystalline texture of BTO and other perovskite ferroelectric films is generally determined by the epitaxial stress imposed by single crystalline perovskite substrates (compressive and tensile favor a-orientation and c-orientation, respectively)1,16-17 as well as the oxygen partial pressure during growth by PLD.18-22 The relevance of the oxygen pressure in the deposition chamber is due to two reasons. On one hand, the oxygen present in the ceramic BTO target using for PLD is insufficient for complete oxidation of the growing film (oxygen has low vapor pressure and it is easily reevaporated). Thus, films deposited under low oxygen pressure can present oxygen vacancies. On the other hand, the laser generated plasma contains high energy ions and atoms, the energy depending on the chamber oxygen pressure.23 Under vacuum conditions, the PLD plasma expands free and atoms and ions having large energy of even tens of eV impact the substrate.24 However, in presence of an oxygen background, the energy of the PLD plasma decreases as higher is the oxygen pressure due to collisions.23,25 The energetic plasma can generate point defects when it impacts on the film,24,26 and thus the amount of defects is higher as lower in the oxygen pressure. Oxygen vacancies and defects, both favoured under low oxygen pressure deposition, cause expansion of the unit cell of the oxide film.12,27-28 The clamping of the thin film to the substrate can favour expansion along the out-of-plane direction, and indeed films deposited at low pressure (more defects) are typically c-oriented and those grown at high pressure (less defects) are a-oriented.18-21 However, to the best of our knowledge, the control of the texture by changing the deposition oxygen pressure during growth has not been achieved so far nor for BTO nor for other ferroelectric oxides epitaxially integrated on silicon wafers. Here we demonstrate texture and strain control in BTO films grown epitaxially on Si(001) buffered with yttria-stabilized zirconia (YSZ), CeO2 and LaNiO3 (LNO). The reason for


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using such a complex heterostructure is as follows. The first layer, YSZ, grows epitaxially on Si(001), the CeO2 and LNO layers are used to accommodate progressively the high lattice mismatch between YSZ and BTO. Moreover, the conducting LNO layer can be used as a bottom electrode. A series of BTO/LNO/CeO2/YSZ/Si(001) samples was prepared varying the oxygen pressure during BTO growth. This parameter can affect the oxygen stoichiometry, and it modulates the kinetic energy of the atoms in the PLD plasma, which can have a severe impact on the structure and defects in the film. We show here that oxygen pressure causes significant changes in BTO tetragonality, and is thus an effective strain engineering tool that can also be used to select the polar axis orientation either in-plane or out-of-plane.

2. EXPERIMENTAL SECTION BTO/LNO/CeO2/YSZ heterostructures were deposited in a single process on Si(001) by PLD (KrF excimer laser, 5 Hz). The BTO films were deposited at 700 °C (temperature measured with a thermocouple inserted in the heater block), and a series of BTO films were prepared under dynamic oxygen pressure (PO2) of 5x10-3, 0.01, 0.015, 0.02, 0.05, and 0.1 mbar. The films were cooled to room temperature under an oxygen pressure of 0.2 mbar. Two additional samples were cooled by switching off the heater immediately and by adding a dwell time of 1 hour at 600 °C. These films were deposited under 0.02 mbar of oxygen and then oxygen partial pressure was increased to 200 mbar of oxygen at the end of the growth. BTO films were deposited with the same number of laser pulses (2000). The growth rate (Figure S1) was in the 0.47-0.56 Å/laser pulse range, being the corresponding thickness range 94-112 nm. Similar series of BTO films were deposited varying oxygen pressure on SrTiO3(001) and LSAT(001) perovskite substrates.


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The lattice parameters were measured by X-ray diffraction (XRD), from symmetrical θ-2θ scans and reciprocal space maps (RSMs) around asymmetrical reflections. Atomic force microscopy (AFM) in dynamic mode was used to characterize the surface morphology. Platinum top electrodes, 20 nm thick and 60 µm x 60 µm in size, were deposited by dc magnetron sputtering through stencil masks. Scanning transmission electron microscopy (STEM) was used for microstructural characterization of the PO2 = 0.01 mbar and 0.1 mbar samples. An aberration corrected Nion UltraSTEM 200, operated at 200 kV and equipped with a 5th order Nion aberration corrector and a JEOL ARM 200CF STEM with a cold field emission source operated at 200 kV and equipped with a CEOS aberration corrector were used. A high angle annular dark field (HAADF) detector collected images of cross sectional specimens were recorded along the Si[1-10] zone axis. In this imaging mode the contrast between atomic columns scales with the square of the atomic number (Z). The polar distortion in BTO unit cells was determined by fitting 2D-Gaussians to locate Ba and Ti atomic column centers. From these positions, the center of each unit cell was defined as the point where the two diagonals connecting Ba positions meet. The polar distortion was defined as the difference between the unit cell center and its Ti column position. Ferroelectric polarization loops and leakage current curves were measured at room temperature in top-top configuration (two BTO capacitors were measured in series, contacting two top Pt electrodes and using the conducting LNO buffer layer as common bottom electrode) using the AixACCT TFAnalyser2000 platform. Ferroelectric polarization loops were obtained at 10 kHz using the dielectric leakage current compensation (DLCC) method.29-30 Leakage current curves were measured with an integration time of 3 s and averaging the current recorded for increasing and decreasing voltage.


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3. RESULTS AND DISCUSSION The symmetrical XRD θ-2θ scan of the PO2 = 0.02 mbar sample is shown in Figure 1a. The diffraction peaks correspond to (00l) reflections of the silicon substrate, buffer layers or BTO film. Zoomed θ-2θ scans (in the 42 - 75° range) of the samples deposited at varied PO2 are presented in Figure 1b. The patterns are vertically shifted for clarity, from bottom to top as PO2 increases from 5x10-3 to 0.1 mbar. The vertical solid and dashed lines mark the position of (002) and (200) reflections of bulk BTO, respectively. The corresponding reflection of the BTO film shifts monotonically from an angle lower than bulk BTO(002) to an angle higher than bulk BTO(200), pointing to the transition from c-oriented BTO to a-oriented BTO as the oxygen pressure during growth increases. The BTO mosaicity, represented by the rocking curve of the specular BTO(002)/(200) reflection, is similar in all films, with a full width at half maximum (FWHM) in the 1.0-1.2° range (Figure S2). It is remarkable that the FWHM of the BTO(002)/(200) specular reflection is similar in all the samples, in the 0.29-0.41° range (Figure S2), indicating homogeneous, but different and selectable, strain state in the series of films. The out-of-plane (oop) and in-plane (ip) lattice parameters were determined from the XRD symmetrical θ-2θ scans (Figure 1a-b) and asymmetrical RSMs around BTO(203) and Si(224) reflections (Figure 1c-h), respectively. The parameters are plotted against PO2 in Figure 2a. The oop lattice parameter (solid triangles) decreases more than 2% with PO2, from 4.072 Å (PO2 = 5x10-3 mbar) to 3.987 Å (PO2 = 0.1 mbar). The BTO unit cell volume is highly enlarged (3.09%) in the PO2 = 5x10-3 mbar film and shrinks monotonically as the pressure increases (0.32% in the PO2 = 0.1 mbar film), suggesting that the expansion (Figure 2b, right axis) is likely caused by


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defects, either those caused by the energetic PLD plasma under low oxygen pressure31 or oxygen vacancies.32 The oop parameter in the two highest PO2 samples is smaller than the c-axis parameter of bulk BTO, and it is even smaller than the bulk a-parameter in the PO2 = 0.1 mbar sample. In contrast, the ip parameter (open squares) of films shows less variation on deposition pressure. The ratio between oop and ip parameters (Figure 2b) in the films deposited at 0.02 mbar and lower pressures is larger than 1 and close to the c/a ratio in bulk BTO (around 1.01), indicating that the BTO cell is elongated along the out-of-plane direction. For the two films deposited at 0.05 and 0.1 mbar, the oop/ip ratio is smaller than 1 and in the PO2 = 0.1 mbar film reaches the a/c ratio in bulk (around 0.99), signalling that these two films are a-oriented. Topographic atomic force microscopy images (Figure S3) show that rms surface roughness of the a-oriented films is notably higher (rms up to around 2.5 nm) than that of the c-oriented films (0.5 – 0.9 nm). The distinct a- or c-orientation of the BTO films depending on the oxygen pressure during deposition is sketched in the inset of Figure 2a. Remarkably, this effectiveness of the oxygen pressure as a tuning knob to tune BTO lattice strain and orientation is not limited to films deposited on Si(001) as a particular substrate. To demonstrate it, two series of BTO films were deposited varying oxygen pressure on bare SrTiO3(001) and LSAT(001) perovskite substrates. AFM (Figure S4-1) XRD (Figure S4-2) characterizations show the same dependence of lattice strain, orientation and surface roughness on pressure. STEM-HAADF cross-sectional images of the PO2 = 0.01 mbar and PO2 = 0.1 mbar samples were recorded for structural characterization at atomic scale (low magnification images of the heterostructures are shown in Figure S5). High magnification Z-contrast images of the BTO films permitted mapping of Ba and Ti column positions (Figure 3). BTO deposited under PO2 =


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0.01 mbar is very homogeneous except for the presence of antiphase boundaries extending vertically across the film (see Figure 3a, and the zoomed area on the right). Antiphase boundaries are likely formed at the LNO/CeO2 interface, as the CeO2 nominal steps height of 5.41 Å is approximately 1.5 times larger than the corresponding height of LNO steps (3.84 Å). The PO2 = 0.1 mbar film presents similar antiphase boundaries, but also contains a high density of voids, having inhomogeneous width of 1-3 nanometers and extending tens of nanometers linearly, close to 15° away to the normal (Figure 3b). Voids are formed by coalescence of three-dimensional BTO islands that grow under high oxygen pressure (Figure S3). The high resolution of the aberration corrected STEM images makes it possible to measure picometer atomic displacements. We have extracted the relative displacement vector maps of Ti with respect to Ba, with the purpose of obtaining information on the polarization configuration in the samples. HAADF original images and with the corresponding Ti displacement map superimposed are shown in Figures 3c (PO2 = 0.01 mbar sample) and 3d (PO2 = 0.1 mbar sample). These dipole maps have been extracted from maps with a larger field of view (see details of the algorithms used and the full dipole maps in the Figure S6). The maps clearly reveal two configurations: one in which all the Ti displacement vectors point mainly in the out-of-plane direction (PO2 = 0.01 mbar film), and other in which the vectors point mainly in the in-plane direction (PO2 = 0.1 mbar film). This is also seen in the average of the displacement vectors at horizontal planes, plotted on the right side of the corresponding images with superimposed dipoles map. Figures 3c and 3d are thus showing that BTO films are composed by a-domains for high pressure deposition and c-domains for low pressure deposition. Figure 4a shows polarization loops of the BTO films deposited at varied oxygen pressure. The films are ferroelectric, with coercive voltage Vc and remnant polarization Pr in the 0.9 – 7.2 V


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(measuring two capacitors in series) and 1.3 - 8.0 µC/cm2 ranges, respectively. Note that near the maximum applied electric field the loops show a sizeable polarization aperture between the increasing and decreasing voltage curves. This aperture is caused by small leakage current,33-34 which contribution to the Pr is small because data have been recorded using the DLCC measurement method. The dependence of Vc and Pr on pressure is plotted in Figure 4b. The high coercive voltage in the two films deposited at low pressure (above 6 V, corresponding to a coercive electric field Ec higher than 250 kV/cm) decreases to around 1 V (Ec ∼ 50 kV/cm) in the PO2 = 0.05 and 0.1 mbar samples. The coercive electric field of the PO2 = 0.02 mbar film fits perfectly in the Ec – thickness scaling dependence reported for epitaxial BTO films on Si(001) deposited under this pressure.35 The dependence of the remnant polarization is akin to the coercive voltage, with low Pr values in the high PO2 films and enhanced polarization for the lower pressure films excluding the PO2 = 5x10-3 mbar one. Thus, ferroelectric polarization and coercive field are controlled by the oxygen pressure. This deposition parameter determines the amount of defects in the BTO film. The effect of the defects is the expansion of the BTO unit cell with increased out-of-plane lattice parameter. There are more defects as lower is the deposition oxygen pressure, causing c-orientation in films deposited below around 0.05 mbar with expanded oop parameter and enhanced polarization at lower pressure. It is to be noted that the PO2 = 5x10-3 mbar film could only be measured at very low frequency (the loop in Fig. 4a was recorded at 150 Hz), most probably due to degradation of the LNO electrode. Indeed, a detailed view of the XRD pattern of the PO2 = 5x10-3 mbar sample (Figure 1b) shows low intensity of the LNO(002) reflection, likely due to the low stability of LNO under high temperature and low pressure.36 Whereas it does not influence the epitaxial growth of BTO, the


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compensation of the BTO polarization charges would be less efficient degrading the ferroelectric properties (Figure S7). The current-voltage curves of the films (Figure 4c) do not show big differences, although leakage in the PO2 = 0.015 and 5x10-3 mbar samples is higher in most of the voltage range. The observation of similar current-voltage curves in BTO films having different microstructure suggests that current leakage through BTO is dominated by the Schottky barriers at the Pt/BTO and BTO/LNO interfaces. The Schottky barriers at the electrode-ferroelectric interfaces are responsible of the strong dependence of leakage on applied electric field.37 The current density at 1 and 5 V is plotted against the oxygen pressure in Fig. 4d. The leakage is similarly low as it is in high quality ferroelectric films on silicon38 and perovskites11,21-22,39 substrates. To evaluate the stability of the defects, and particularly those involving oxygen vacancies, two films were cooled under high oxygen pressure (200 mbar), one of them with additional annealing for 1 hour at 600 °C (XRD θ-2θ scan and RSMs are presented in Figure S8). The lattice parameters of the annealed samples (Figure 2a) are similar to those of the sample cooled under 0.2 mbar. Moreover, the annealed samples present polarization loops, remnant polarization (Figure 4b) and leakage (Figure 4d) very similar to the corresponding sample cooled down under standard conditions (0.2 mbar). The results signal that low oxygen vacancy concentration and/or highly stable defects ultimately determine lattice strain in the BTO films deposited by PLD, and that their dependence on the deposition oxygen pressure permits growth of films with selectable strain and polar orientation.

3. CONCLUSIONS


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In conclusion, the oxygen pressure during pulsed laser deposition can be used as a tuning knob to control lattice strain of epitaxial BTO over a broad range, and concomitantly ferroelectric polarization without important variations on leakage current. It also permits selecting the polar axis to be in-plane (a-oriented BTO) or out-of-plane (c-oriented BTO). The unconventional strain engineering is efficient for films as thick as 100 nm, and it does not require selection of a specific substrate. Remarkably, the control of polar axis orientation is done for BTO films on either perovskite substrates or Si(001) wafers. Important limitations of conventional strain methodology are thus overcome, opening new avenues for better control of strain and polar orientation in ferroelectric oxides and particularly for its integration on semiconducting substrates.


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BTO(003)/(300) YSZ(004)/LNO(003)


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Figure 1. (a) Symmetrical XRD θ-2θ scan of the PO2 = 0.02 mbar sample. (b) Zoomed θ-2θ scans of the samples deposited at varied PO2. Diffractograms are shifted vertically for clarity. The vertical solid and dashed lines mark the position of the (002) and (200) reflections in bulk BTO, respectively. (c-h) XRD reciprocal space maps around Si(224) and BTO(203) asymmetrical reflections.


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Figure 2. (a) In-plane (open squares) and out-of-plane (solid triangles) lattice parameters as a function of deposition pressure. The horizontal dashed lines mark the a and c lattice parameters in bulk BTO. Circles and rhombi correspond to samples that were cooled under 200 mbar of oxygen and with additional annealing during 1 h at 600 °C, respectively. Inset: Sketch of the BTO orientation depending on the deposition pressure. (b) Tetragonality and unit cell volume as a function of deposition pressure. Horizontal dashed lines mark c/a, a/c and unit cell volume in bulk BTO.


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Figure 3. STEM-HAADF cross-sectional images along the BTO[010] zone axis of BTO films. Antiphase boundary extending vertically in the BTO film deposited under PO2 = 0.01 mbar (a). Defective regions with locally incoherent grain boundaries in the PO2 = 0.1 mbar sample (b). HAADF image together with their corresponding Ti displacement vectors superimposed on the images for the 0.01mbar (c) and 0.1mbar (d) samples. The sketches in the HAADF images illustrate the distinct Ti displacement in each sample. The average of the vectors along horizontal planes is plotted at the right side of the corresponding images with superimposed Ti displacement map.


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Figure 4. (a) Polarization loops of BTO films deposited at varied oxygen pressure on buffered Si(001). (b) Dependence of remnant polarization (left axis, solid symbols) and coercive voltage (right axis, open symbols) on the deposition pressure. Circles (1) and rhombi (2) correspond to the samples cooled down under 200 mbar and with additional annealing during 1 h at 600 °C under 200 mbar, respectively. Leakage curves of BTO films deposited at varied oxygen pressure (c) and leakage current at 1 V (red squares) and 5 V (blue triangles) as a function of the deposition pressure. Black symbols indicate the leakage current at 1 and 5 V of the two annealed films.


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ASSOCIATED CONTENT Supporting Information. Growth rate of BTO as a function of the deposition oxygen pressure. Dependence with the pressure of the width of specular reflections and rocking curves of BTO films on Si(001). Topographic AFM images and roughness of the BTO films on Si(001). XRD and AFM of two equivalent series of BTO films grown on SrTiO3(001) and LSAT(001) perovskite substrates. Low magnification cross-sectional STEM images. Dipole maps with large field of view and description of the mathematical procedure used. Effects of degradation of LNO bottom electrode on the polarization loops. Influence of high oxygen pressure during cooling down and annealing on the lattice strain and polarization loops of BTO films on Si(001).

AUTHOR INFORMATION Corresponding Author *Email: [email protected] ORCID Florencio Sánchez: 0000-0002-5314-453X Notes The authors declare no competing financial interest.

Author Contributions


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The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT Financial support from the Spanish Ministry of Economy and Competitiveness, through the “Severo Ochoa” Programme for Centres of Excellence in R&D (SEV-2015-0496) and the MAT2017-85232-R (AEI/FEDER,UE), MAT2014-56063-C2-1-R, and MAT2015-73839-JIN projects, and from Generalitat de Catalunya (2014 SGR 734) is acknowledged. JL is financially supported by China Scholarship Council (CSC) with No. 201506080019. JL work has been done as a part of his Ph.D. program in Materials Science at Universitat Autònoma de Barcelona. J.G. acknowledges the RyC contract (RC-2012–11709). Electron microscopy observations at ORNL were supported by the U.S. Department of Energy (DOE), Basic Energy Sciences (BES), Materials Sciences and Engineering Division. Authors acknowledge the ICTS-CNME for offering access to their instruments and expertise. S.E. acknowledges the Spanish Ministry of Economy and Competitiveness for his PhD contract (SEV-2015-0496-16-3).

REFERENCES (1) Schlom, D.G.; Chen, L.Q.; Eom, C.B.; Rabe, K.M.; Streiffer, S.K.; Triscone, J.M. Strain Tuning of Ferroelectric Thin Films. Annu. Rev. Mater. Res. 2007, 37, 589-626.


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Control of Polar Orientation and Lattice Strain in Epitaxial BaTiO

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