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A Strain Driven Antiferroelectric-to-Ferroelectric Phase Transition in La-Doped BiFeO3 Thin Films on Si Deyang Chen, Christopher T. Nelson, Xiaohong Zhu, Claudy R. Serrao, James D. Clarkson, Zhe Wang, Ya Gao, Shang-Lin Hsu, Liv R. Dedon, Zuhuang Chen, Di Yi, Heng-Jui Liu, Dechang Zeng, Ying-Hao Chu, Jian Liu, Darrell Schlom, and Ramamoorthy Ramesh Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.7b03030 • Publication Date (Web): 16 Aug 2017 Downloaded from http://pubs.acs.org on August 17, 2017
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A Strain Driven Antiferroelectric-to-Ferroelectric Phase Transition in LaDoped BiFeO3 Thin Films on Si Deyang Chen,*,†,‡,§ Christopher T. Nelson,† Xiaohong Zhu,† Claudy R. Serrao,† James D. Clarkson,† Zhe Wang,∥ Ya Gao,† Shang-Lin Hsu,† Liv R. Dedon,† Zuhuang Chen,† Di Yi,† Heng-Jui Liu,⊥ Dechang Zeng,§ Ying-Hao Chu,⊥ Jian Liu,# Darrell G. Schlom, ∥
,∇
and
Ramamoorthy Ramesh†,¶,● †
Department of Materials Science and Engineering, University of California, Berkeley, Berkeley, California 94720, USA
‡
Institute for Advanced Materials and Guangdong Provincial Key Laboratory of Quantum Engineering and Quantum Materials, South China Normal University, Guangzhou 510006, China §
School of Materials Science and Engineering, South China University of Technology, Guangzhou 510640, China
∥
Department of Materials Science and Engineering, Cornell University, Ithaca, New York 14853, USA ⊥
Department of Materials Science and Engineering, National Chiao Tung University, Hsinchu 30010, Taiwan #
Department of Physics Tennessee 37996, USA
∇Kavli ¶
and
Astronomy,
University
of
Tennessee,
Knoxville,
Institute at Cornell for Nanoscale Science, Ithaca, New York 14853, USA
Department of Physics, University of California, Berkeley, Berkeley, California 94720, USA
●
Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, USA
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ABSTRACT A strain driven orthorhombic (O) to rhombohedral (R) phase transition is reported in Ladoped BiFeO3 thin films on silicon substrates. Biaxial compressive epitaxial strain is found to stabilize the rhombohedral phase at La concentrations beyond the morphotropic phase boundary (MPB). By tailoring the residual strain with film thickness we demonstrate a mixed O/R phase structure consisting of O phase domains measuring tens of nanometers wide within a predominant R phase matrix. A combination of piezoresponse force microscopy (PFM), transmission electron microscopy (TEM), polarization-electric field hysteresis loop (P-E loop) and polarization maps reveal that the O-R structural change is an antiferroelectric to ferroelectric (AFE-FE) phase transition. Using scanning transmission electron microscopy (STEM), an atomically sharp O/R MPB is observed. Moreover, x-ray absorption spectra (XAS) and x-ray linear dichroism (XLD) measurements reveal a change in the antiferromagnetic axis orientation from out of plane (R-phase) to in plane (O-phase). These findings provide direct evidence of spin-charge-lattice coupling in La-doped BiFeO3 thin films. Furthermore, this study opens a new pathway to drive the AFE-FE O-R phase transition and provides a route to study the O/R MPB in these films.
KEYWORDS BiFeO3, multiferroic, antiferroelectric, strain engineering, spin-charge-lattice coupling, antiferromagnetic
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Epitaxial strain is a powerful pathway to induce phase transitions in oxide thin films,1-4 e.g., the strain induced ferroelectric to ferroelectric (FE-FE) phase transition from rhombohedrallike (R) phase to tetragonal-like (T) phase in BiFeO3 films, with enhanced piezoelectricity near the R/T morphotropic phase boundary (MPB) and enhanced magnetism in the constrained R phase between T phases.5-7 Much effort has been focused on the strain driven R to T phase and R/T MPB.8-11 Recently, theoretical calculations for BiFeO3 suggested a straintemperature phase diagram showing different symmetries, including tetragonal, distorted rhombohedral and orthorhombic (O) phases.5,12,13 Meanwhile, the R to O phase transition induced by hydrostatic pressure14,18 or chemical substitution14-19 has attracted much interest due to the study of competing spin, charge, orbital, and lattice degrees of freedom as well as the possible presence of a MPB. Here, we study the strain engineering of the O to R phase transition and the possible formation of an O/R MPB. Based on our prior work on the formation of a R/T MPB in the BiFeO3 model system,5-7 we are prompted to explore three questions: (i) is it possible to achieve a strain driven O to R phase transition and does the O/R MPB exist in this system, (ii) is the O to R phase transition also a FE-FE phase transition, and (iii) how does the phase transition affect the magnetic ordering? In this study, we start from films with a nominal composition of La0.22Bi0.78FeO3 (LBFO), which possess an equilibrium orthorhombic crystal structure, but is very close to the rhombohedral-orthorhombic phase boundary.19 Chemical analysis of the films was performed via Rutherford backscattering spectrometry (RBS), as shown in Figure S1. We have discovered a strain driven antiferroelectric to ferroelectric (AFE-FE) phase transition from orthorhombic (O) to rhombohedral (R) phase by growing these films on SrTiO3-buffered Si substrates prepared by reactive molecular-beam epitaxy (MBE).20-22 When the LBFO films are thin (20 nm) they experience biaxial compressive epitaxial strain to the SrTiO3-buffered Si substrates on which they are grown, driving them into the R phase. Thick (125 nm) LBFO films undergo full strain relaxation and consequently are O phase. In the O/R mixed phase 3 ACS Paragon Plus Environment
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region, we show that O and R phases can coexist and an atomically sharp O/R morphotropic phase boundary (MPB) is observed. Finally, x-ray absorption spectra (XAS) and x-ray linear dichroism (XLD) measurements reveal that the antiferromagnetic axis orientation of LBFO is changed from out-of-plane (R-phase) to in-plane (O-phase). We studied a series of epitaxial LBFO films with thickness of 20-125 nm grown on 30 nm thick SrRuO3 (SRO) bottom electrode layers on top of (001) Si substrates buffered by 20 nm thick SrTiO3 (STO) layers, where the STO was deposited by MBE. The LBFO and SRO were grown by pulsed laser deposition (PLD) at 690 °C in an oxygen pressure of 6 ×10-2 Torr and cooled in a 1 atm oxygen atmosphere. A combination of x-ray diffraction (XRD), reciprocal space mapping (RSM) and transmission electron microscopy (TEM) was used for the structural characterization. Ferroelectric properties were measured using piezoresponse force microscopy (PFM) and ferroelectric tester. X-ray absorption spectra (XAS) and x-ray linear dichroism (XLD) measurements were performed at the beamline 4.0.2 of the Advanced Light Source, Lawrence Berkeley National Laboratory, to study the magnetic ordering of LBFO films. The total electron yield detection mode was used. The grazing angle of the incoming beam was fixed at 30 degrees while the x-ray polarization was rotated to obtain XLD. A cross sectional image of an LBFO/SRO/STO/Si sample is shown in Figure 1a taken by high-angle annular dark field (HAADF) scanning-TEM (STEM). This type of image is sensitive to variations in atomic number, and the composition changes between the layers of the heterostructure are clearly distinguished. The image shows that the samples have high quality epitaxial growth, with sharp interfaces. A high-resolution Z-contrast image of the LBFO/SRO interface (Figure 1b) illustrates the atomic-scale epitaxy between the SrRuO3 and LBFO layers. High quality epitaxial growth of LBFO films is further confirmed by atomic force microscope (AFM) , as presented in Figure S2, showing an atomically flat morphology with a root mean square (RMS) roughness of only one unit cell (0.4 nm).
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Let us now turn to our first question: is it possible to achieve a strain driven O to R phase transition? Figure 1c shows a typical XRD θ−2θ scan of LBFO thin films with thicknesses 20, 35, 60, 80, and 125 nm, which further confirms the epitaxial growth on Si. We observe that the 125 nm film is the orthorhombic phase, which is in agreement with an earlier study;19 detailed RSM scans in Figure S3 show that the experimental data agree well with the orthorhombic structure, confirming that it is the O-phase, which yields only two peaks occurring along each of the pseudocubic H0L and HHL directions, while the typical monoclinic structure of BFO yields 2 peaks along the H0L and 3 peaks along the pseudocubic HHL. As shown in Figure 1c, the fraction of O phase decreases and R phase increases when the film thickness is reduced to 85 nm and 60 nm, manifested by the relative intensity of the two peaks, but the relative intensity is almost the same in the 35 nm film whereas the second peak is even slightly stronger than the 002 peak in the 20 nm film. We expect that the 002 peaks of 30 nm SRO bottom electrode and 20 nm STO buffer layer contribute dominantly to the second peak in the 20 nm and 35 nm thick LBFO films. Thus, we infer that a strain driven O to R phase transition exists in the LBFO films due to biaxial compressive epitaxial strain from the STO-buffered Si substrates as the film thickness is reduced from 125 nm to 20 nm. To assess this hypothesis, we grew LBFO films with a range of thicknesses (20-125 nm) on STO-buffered Si substrates without SRO. Figure 1d shows typical θ−2θ scans of 20 nm, 35 nm, and 125 nm thick films. Also, the XRD scan of the STO-buffered Si substrate was used as the reference. As expected, the high angle peak of a 20 nm thick film corresponds precisely to the STO 002 peak of the STO-buffered Si substrate, which demonstrates that the 20 nm film is pure R phase. In contrast, the 35 nm film is mixed O and R phase with only a small amount of O phase, while the 125 nm film is the O-phase. In addition, the peak shift with decreasing thickness shows the biaxial compressive strain effect from the STO-buffered Si substrates. We can therefore conclude that a strain driven O- to R-phase transition is achieved as a function of thickness. It is important to note that the thermal mismatch between 5 ACS Paragon Plus Environment
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Si (αSi ~ 3 ×10−6 K−1) and the oxides (αBFO ~ 1 × 10−5 K−1) introduces a tensile thermal strain which favors the O phase in thick films (The XRD data of 125 nm thick LBFO on STO substrate and Si substrate, as presented in Figure S4, show that a tensile strain is introduced due to the thermal expansion effect), while the compressive strain imposed by STO-buffered Si plays a predominant role in thin films (Figure 1d), favoring the R phase. We probe the piezoelectric properties of the LBFO samples as a function of film thickness (20 nm~125 nm) using piezoresponse force microscopy (PFM). Figure 2 shows the out-ofplane (OOP) and in-plane (IP) electrical switching behavior of the samples. 5 µm × 5 µm OOP and IP PFM images of all of the samples were taken after poling 3 µm × 3 µm boxes at various DC field from -8 V to -15 V. First, we noticed that the as-grown domain size is much smaller in LBFO films compared to pure BiFeO3 films.23 This significant difference in domain structure results from the La doping in the LBFO films.19 Distinct switching behavior changes in OOP contrast and IP contrast are observed in the films with different thicknesses. The OOP and IP PFM images of O phase 125 nm thick film shows a relatively negligible contrast after -15 V DC field poling, implying the O phase is non-ferroelectric (Figure 2a). According to the OOP and IP contrast in PFM images, however, we observe that the switching of the polarization becomes progressively stronger with a reduction of film thickness, from 80 nm to 35 nm (Figure 2b and c), suggesting enhanced ferroelectricity with the increase of R phase. By applying an ac modulation and dc bias simultaneously on the PFM tip, local piezoresponse phase loops were taken on 35 nm, 80 nm, and 125 nm thick LBFO films, respectively. Figure 2d shows 180° sharp ferroelectric switching of the 35 nm thick film (composed primarily of R phase) and weak switching of 80 nm thick film (O and R mixed phases) whereas no switching was observed in the O phase 125 nm thick film. Therefore, the PFM data (Figure 2a) and the red phase loop measurement (Figure 2d) prove that the O phase 125 nm thick film is non-ferroelectric, while the related characterizations as shown in Figure 6 ACS Paragon Plus Environment
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2c and Figure 2d (the black phase loop) demonstrate that the 35 nm thick film with principal R phase is ferroelectric. Further P-E loop study as shown in Figure 2e reveals the antiferroelectric double hysteresis loop of O phase 125 nm thick film (the expected ferroelectric single hysteresis loop was not obtained due to the leak problem in the 35 nm thin film). These results indicate that the O-R structural change, with the decrease of film thickness, is antiferroelectric to ferroelectric (AFE-FE) phase transition. To understand the fine structure of the orthorhombic and rhombohedral mixed phase, we now focus on the details of a 35 nm thick film. We carried out further TEM and STEM-based imaging studies. Figure 3a shows the cross-sectional dark field diffraction contrast TEM image of the 35 nm film that exhibits a mixture of O and R phases. In this image, there are two O phase domains, which have adopted mirrored orientations of ordered superlattice structures along (101) planes. From selected area electron diffraction (SAED) patterns, Figure 3b, this superlattice ordering corresponds to the highlighted 1/4 (101) reflections indicating a quadrupling of the pseudocubic unit cell with the periodicity ~ 11.3 Å. Each of the two Odomain orientations in Figure 3a contribute unique superlattice reflections in the SAED image. The spots in red circles in Figure 3b arise from region I which is one of the orientations of the orthorhombic phase as shown in Figure 3a; the spots in the blue squares are from the other orientation in region II. The diffraction pattern we report here was also observed in Sm-doped BiFeO3 films,17,24 which is an antiferroelectric (AFE) orthorhombic PbZrO3 (PZO)-like structure.25,26 This again points that the O phase LBFO film is antiferroelectric, consistent with the double hysteresis loop measurement in Figure 2e. It is also important to note that the dark field large area TEM image in Figure 3c shows that the film is predominantly R phase and only a small fraction of the O phase exists in a 35 nm thick film (the arrows indicate that the bright areas are the O phase), which is in agreement with the XRD data in Figure 1. Atomic scale HAADF STEM imaging was performed to reveal the structural details of the mixed O and R phase. Figure 4a shows a boundary between the O and R phases depicted with 7 ACS Paragon Plus Environment
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a dashed white line. The transition between O and R phases is easily visible from a colorized map of the A-site atomic spacing along the in-plane [100] axis in Figure 4b, depicting the Asite atomic spacing calculated from a Gaussian fit to the atomic positions. The orthorhombic superlattice structure is clearly manifested in the A-site strain map, highlighting the O region in the upper left portion of the image. The O/R phase boundary aligns to the (101) plane and several steps are visible, indicated by the dashed step lines in Figure 4b. The transition between O and R phases is without dislocations or defects and appears nearly atomically sharp. The FFT obtained from a full image of Figure 4a, an R phase region, and an O Phase region are shown in Figure S5. The FFT of the R phase shows the same symmetry as the R phase in pure BiFeO3,27-29 while the FFT of the O phase is consistent with the SAED patterns in Figure 3b, exhibiting 1/4 (101) superlattice reflections. As a result, the SAED patterns in Figure 3b and FFT in Figure S5 reveal clear differences between the symmetries of the O and R phases. The fine structure analysis of the O phase from atomic scale STEM corroborates the antiferroelectric structure indicated by PFM and P-E loop measurements (Figure 2). We now focus on the detailed structure of mixed O and R phase in the 35 nm thick film to assess the space group of the O phase. Figure 4a and b shows that the upper left orthorhombic phase consists of alternating diagonal shifts of the La/Bi site whose spacing is consistent with the observed diffraction spots in Figure 3b, while the lower right rhombohedral phase has uniform spacing. These results again indicate that the O-phase is antiferroelectric (AFE) whereas the R-phase is ferroelectric (FE). Besides, the cation positions of La/Bi, Fe and O have been observed in the annular bright field (ABF) STEM image of O phase region, as shown in Figure 4c. The corresponding fitting image of Figure 4c, as shown in Figure 4d, demonstrates that the projected structure closely matches PbZrO3 (atomic model shown overlaid). In this case it is a centrosymmetric Pbam unit cell. Thereby, the cation positions and the symmetry observed in the SAED patterns reveal that the O phase in LBFO film is an 8 ACS Paragon Plus Environment
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antiferroelectric orthorhombic Pbam structure. Furthermore, the polarization mapping defined by cation displacement of a cross section of the O/R interface are shown in Figure 4e, which has been averaged over 5 adjacent cells to reduce image noise. A yellow arrow is depicted on each of the Bi/La and Fe atoms corresponding to their offset from the average of their four nearest neighbors (inverted for B sites). Within the rhombohedral region this offset is uniform along [101] corresponding to the known long-range non-centrosymetric [111] displacement of the A-site which gives the R phase a large net polarization.30,31 The O phase, in contrast, shows displacements along the [101] which alternate in sign, corresponding to antidistortive oscillation of the A-site along the observed superlattice ordered direction. This again strongly points to an antiferroelectric (AFE) phase at the atomic level. We can therefore conclude that the O-R phase transition in LBFO films is AFE-FE, which demonstrates the charge-lattice coupling in LBFO films. To explore the magnetic ordering and the spin-lattice coupling, we investigated the magnetism of a pure BFO film that is entirely R phase (as a reference sample), a 10 nm thick LBFO film that is also R phase, and 125 nm thick LBFO film that is O phase by carrying out Fe L-edge x-ray absorption spectra (XAS) and x-ray linear dichroism (XLD) measurements. Figure 5a shows the Fe L2,3-edge XAS of the 125 nm thick O phase LBFO film with the x-ray polarization perpendicular (E⊥c) and parallel (E//c) to the film surface. XLD data of the pure BFO, 10 nm, and 125 nm LBFO films were obtained by subtracting the E//c spectrum from the E ⊥ c one, as shown in Figure 5b. The result reveals that the orientation of antiferromagnetic axis of the 10 nm R phase LBFO film (black line) is consistent with the pure BFO reference sample with R-phase (blue line), which is in agreement with previous studies.32, 33 This further confirms that the R phase is induced from the O phase by the strain, which is, once again, consistent with the XRD and TEM data. More interestingly, the sign of the XLD spectrum is reversed between the R phase and the O phase samples, which can be 9 ACS Paragon Plus Environment
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most clearly seen at the Fe L2-edge. It is noteworthy that the XLD spectrum could originate from the antiferromagnetic state and (or) the FE/AFE state of the films, since the antiferroic polarization of the O-phase is pointing along [101] while the ferroic polarization of the Rphase is along [111] (demonstrated by the TEM in Figure 4). That means the polarizationinduced distortion to the FeO6 cage would have a similar in-plane and out-of-plane components for both phases (the R-phase is slightly more in-plane). Thus, one would expect the contribution to XLD of both R and O phases to be similar or to at least have the same sign. The XLD data, however, show a sign switch. So we conclude that the antiferromagnetic axis orientation is the main driving force. Therefore, this observation of the reversed XLD spectrum indicates that the antiferromagnetic axis orientation is changed from primarily out of plane in a 10 nm thick R-phase film to primarily in plane in a 125 nm thick O phase LBFO film (red line). These findings demonstrate the spin-lattice coupling in LBFO thin films. In summary, our studies have revealed the ability of a biaxial compressive strain induced orthorhombic to rhombohedral phase transition in La-doped BiFeO3 thin films on Si. Furthermore, we find the existence of an O/R mixed phase with an atomically sharp O/R morphotropic phase boundary. Combined with PFM, P-E hysteresis loop measurement and high resolution TEM imaging, we demonstrate that the orthorhombic to rhombohedral phase transition is an antiferroelectric to ferroelectric phase transition. XAS and XLD measurements enable the discovery of the change of the antiferromagnetic axis orientation from out of plane (R phase) to in plane (O phase). Our findings reveal the direct evidence of the spin-chargelattice coupling in La-doped BiFeO3 thin films. Moreover, this study opens a new pathway to drive the AFE-FE O-R phase transition and provides a route to study the O/R MPB in these films. Supporting Information Rutherford backscattering spectrometry (RBS) measurement, atomic force microscope (AFM) images of LBFO films, reciprocal space mapping (RSM) results on the 125 nm thick LBFO 10 ACS Paragon Plus Environment
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film, x-ray diffraction (XRD) patterns of 125 nm thick LBFO films grown on STO and STO buffered Si, and the Fast Fourier transform (FFT) data obtained from O/R mixed phase, R phase region, and O Phase region.
AUTHOR INFORMATION Corresponding Authors *E-mail:
[email protected] Author Contributions D. Y. Chen, C. T. Nelson and X. H. Zhu contributed equally to this work.
Notes The authors declare no competing financial interest.
Acknowledgements This work was primarily funded by the National Science Foundation (Nanosystems Engineering Research Center for Translational Applications of Nanoscale Multiferroic Systems) under grant number EEC-1160504. Electron microscopy was performed at National Center for Electron Microscopy (NCEM), LBNL which is supported by the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DEAC02-05CH11231. Use of the Advanced Light Source is supported by the Director, Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. D.Y. C. acknowledges financial support from the Oversea Study Program of Guangzhou Elite Project (GEP). J.L. acknowledges financial support by the Science Alliance Joint Directed Research and Development Program at the University of Tennessee.
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(12) Yang, J. C.; He, Q.; Suresha, S. J.; Kuo, C. Y.; Peng, C. Y.; Haislmaier, R. C.; Motyka, M. A.; Sheng, G.; Adamo, C.; Lin, H. J.; Hu, Z.; Chang, L.; Tjeng, L. H.; Arenholz, E.; Podraza, N. J.; Bernhagen, M.; Uecker, R.; Schlom, D. G.; Gopalan, V.; Chen, L. Q.; Chen, C. T.; Ramesh, R.; Chu, Y. H. Phys. Rev. Lett. 2012, 109, 247606. (13) Infante, I. C.; Lisenkov, S.; Dupé, B.; Bibes, M.; Fusil, S.; Jacquet, E.; Geneste, G.; Petit, S.; Courtial, A.; Juraszek, J.; Bellaiche, L.; Barthélémy, A.; Dkhil, B. Phys. Rev. Lett. 2010, 105, 057601. (14) Knee, C. S.; Tucker, M. G.; Manuel, P.; Cai, S.; Bielecki, J.; Börjesson, L.; Eriksson, S. G. Chem. Mater. 2013, 26, 1180-1186. (15) Borisevich, A. Y.; Eliseev, E. A.; Morozovska, A. N.; Cheng, C. J.; Lin, J. Y.; Chu, Y. H.; Kan, D.; Takeuchi, I.; Nagarajan, V.; Kalinin, S. V. Nat. Commun. 2012, 3, 775. (16) Kan, D.; Cheng, C.-J.; Nagarajan, V.; Takeuchi, I. J. Appl. Phys. 2011, 110, 014106. (17) Cheng, C. J.; Kan, D.; Lim, S. H.; McKenzie, W. R.; Munroe, P. R.; Salamanca-Riba, L. G.; Withers, R. L.; Takeuchi, I.; Nagarajan, V. Phys. Rev. B 2009, 80, 014109. (18) Catalan, G.; Scott, J. F. Adv. Mater. 2009, 21, 2463-2485. (19) Chu, Y. H.; Zhan, Q.; Yang, C. H.; Cruz, M. P.; Martin, L. W.; Zhao, T.; Yu, P.; Ramesh, R.; Joseph, P. T.; Lin, I. N.; Tian, W.; Schlom, D. G. Appl. Phys. Lett. 2008, 92, 102909. (20) Mi, S. B.; Jia, C. L.; Vaithyanathan, V.; Houben, L.; Schubert, J.; Schlom, D. G.; Urban, K. Appl. Phys. Lett. 2008, 93, 101913. (21) Baek, S. H.; Eom, C. B. Acta Mater. 2013, 61, 2734-2750. (22) Baek, S. H.; Park, J.; Kim, D. M.; Aksyuk, V. A.; Das, R. R.; Bu, S. D.; Felker, D. A.; Lettieri, J.; Vaithyanathan, V.; Bharadwaja, S. S. N.; Bassiri-Gharb, N.; Chen, Y. B.; Sun, H. P.; Folkman, C. M.; Jang, H. W.; Kreft, D. J.; Streiffer, S. K.; Ramesh, R.; Pan, X. Q.; Trolier-McKinstry, S.; Schlom, D. G.; Rzchowski, M. S.; Blick, R. H.; Eom, C. B. Science 2011, 334, 958-961. 13 ACS Paragon Plus Environment
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(23) Martin, L. W.; Chu, Y.-H.; Zhan, Q.; Ramesh, R.; Han, S.-J.; Wang, S. X.; Warusawithana, M.; Schlom, D. G. Appl. Phys. Lett. 2007, 91, 172513. (24) Kan, D.; Pálová, L.; Anbusathaiah, V.; Cheng, C. J.; Fujino, S.; Nagarajan, V.; Rabe, K. M.; Takeuchi, I. Adv. Funct. Mater 2010, 20, 1108-1115. (25) Viehland, D. Phys. Rev. B 1995, 52, 778. (26) Woodward, D. I.; Knudsen, J.; Reaney, I. M. Phys. Rev. B 2005, 72, 104110. (27) Wang, J.; Neaton, J. B.; Zheng, H.; Nagarajan, V.; Ogale, S. B.; Liu, B.; Viehland, D.; Vaithyanathan, V.; Schlom, D. G.; Waghmare, U. V.; Spaldin, N. A.; Rabe, K. M.; Wuttig, M.; Ramesh, R. Science 2003, 299, 1719-1722. (28) Das, R. R.; Kim, D. M.; Baek, S. H.; Eom, C. B.; Zavaliche, F.; Yang, S. Y.; Ramesh, R.; Chen, Y. B.; Pan, X. Q.; Ke, X.; Rzchowski, M. S.; Streiffer, S. K. Appl. Phys. Lett. 2006, 88, 242904. (29) Kim, Y. H.; Bhatnagar, A.; Pippel, E.; Alexe, M.; Hesse, D. J. Appl. Phys. 2014, 115, 043526. (30) Nelson, C. T.; Winchester, B.; Zhang, Y.; Kim, S.-J.; Melville, A.; Adamo, C.; Folkman, C. M.; Baek, S.-H.; Eom, C.-B.; Schlom, D. G.; Chen, L.-Q.; Pan, X. Nano Lett. 2011, 11, 828-834. (31) Li, L.; Gao, P.; Nelson, C. T.; Jokisaari, J. R.; Zhang, Y.; Kim, S.-J.; Melville, A.; Adamo, C.; Schlom, D. G.; Pan, X. Nano Lett. 2013, 13, 5218-5223. (32) Zhao, T.; Scholl, A.; Zavaliche, F.; Lee, K.; Barry, M.; Doran, A.; Cruz, M. P.; Chu, Y. H.; Ederer, C.; Spaldin, N. A.; Das, R. R.; Kim, D. M.; Baek, S. H.; Eom, C. B.; Ramesh, R. Nat. Mater. 2006, 5, 823-829. (33) Ko, K.-T.; Jung, M. H.; He, Q.; Lee, J. H.; Woo, C. S.; Chu, K.; Seidel, J.; Jeon, B.-G.; Oh, Y. S.; Kim, K. H.; Liang, W.-I.; Chen, H.-J.; Chu, Y.-H.; Jeong, Y. H.; Ramesh, R.; Park, J.-H.; Yang, C.-H.; Nat. Commun. 2011, 2, 567.
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(b)
(a)
LBFO
LBFO SRO STO
SRO 20 nm
Si
(d)
(c)
O Phase
O+R Phase
O
R Phase
STO(002)
O+R R
Figure 1. (a) HAADF low-resolution STEM image of 35 nm thick LBFO film on STObuffered Si with SRO as bottom electrode. (b) Dark field high-resolution TEM image of LBFO and SRO interface. (c)Typical x-ray θ−2θ scan of LBFO films with different thicknesses (20 nm to 125 nm) on STO-buffered Si with SRO bottom electrodes. (d) θ−2θ scan of STO-buffered Si substrate and LBFO films on STO-buffered Si without an SRO layer.
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Figure 2. OOP PFM (left) and IP PFM (right) images of LBFO films with thicknesses (a) 125 nm, (b) 80 nm, and (c) 35 nm poled with various DC voltages. The black boxes demark the areas that were switched. The scale bars are 1 µm. (d) The piezoresponse phase curves of 35 nm, 80 nm, and 125 nm thick LBFO films. (e) P-E loop of 125 nm thick LBFO film.
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(a)
vacuum
(b)
Orthorhombic phase
glue
vacuum LBFO
LBFO
(001)
(101)
Si
I
II
(100) 500 nm Film relaxed (split peak)
Rhombohedral phase SRO
(c) 10 nm
Figure 3. (a) Dark field cross-sectional TEM image of a mixed O and R phase region in the 35 nm thick LBFO film. (b) Selected area electron diffraction patterns of the mixed O and R phase region. (c) Dark field large area cross-sectional TEM image of the 35 nm thick LBFO film on SRO on STO-buffered Si.
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Figure 4. (a) High angle angular dark field STEM image of the O-R morphotropic phase boundary in the 35 nm thick LBFO film. An atomically sharp O/R phase boundary along (101) planes is observed; no dislocations or defects are found at the interface. (b) A-site in-plane spacing strain map corresponds to the O/R MPB STEM image in (a). (c) Annular bright field STEM image of O phase region in the LBFO film. (d) Inverted bright field fitting image of (c) overlaid with PbZrO3 atomic model. (e) Polarization mapping defined by cation displacement of the O and R mixed phases in the LBFO films. The cross section fine structure of the O-R phase boundary is fitted from the average over 5 adjacent unit cells of the square region denoted by the dashed line as shown in (a). 18 ACS Paragon Plus Environment
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(b)
(a)
Figure 5. (a) Fe L2,3-edge x-ray absorption spectra (XAS) of the 125 nm thick LBFO film consisting of the O phase with the x-ray linear polarization E parallel to c (E//c) and perpendicular to c (E⊥c). (b) Fe L2,3-edge x-ray linear dichroism (XLD) of a BiFeO3 film that is pure R phase, a ~ 10 nm LBFO on Si that is pure R phase, and a ~ 125 nm thick LBFO film on Si that is O phase.
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