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Room-Temperature Ferroelectricity in Hexagonal TbMnO3 Thin Films Dong Jik Kim,* Tula R. Paudel, Haidong Lu, John D. Burton, John G. Connell, Evgeny Y. Tsymbal, S. S. Ambrose Seo, and Alexei Gruverman* Rare-earth manganite materials (ReMnO3) exhibit a wealth of fundamentally exciting and technologically appealing properties. For example, a coupled magnetic and ferroelectric ordering in these materials allows the manipulation of ferroelectric polarization with a magnetic field and/or magnetization with an electric field, thereby opening new possibilities for development of multifunctional electronic devices. A fundamentally intriguing aspect is that, depending on the size of the Re ion, these materials exhibit either an orthorhombic or a hexagonal structure. While the orthorhombic rare-earth manganites, such as TbMnO3, demonstrate ferroelectric ordering at rather low temperatures (well below 100 K), the hexagonal ReMnO3 tend to be high-temperature ferroelectrics with a transition temperature typically above 600 K. Here, we report room-temperature ferroelectricity in thin films of epitaxially stabilized hexagonal TbMnO3, which in the bulk form exhibits orthorhombic structure. Our theoretical modeling reveals that the ferroelectric polarization of hexagonal TbMnO3 arises from the rotation and tilt of the MnO5 polyhedra, similar to other hexagonal manganites. Switchable ferroelectricity is accompanied by significant polarization-dependent resistive switching, which is attributed to polarization-dependent Schottky barrier modulation. The obtained results demonstrate that strain-engineering of rareearth manganites is a viable route to realize new physical properties of these materials and achieve novel functionalities at elevated temperatures useful for practical applications. Rare-earth manganites (ReMnO3) have attracted much attention since the discovery of coupled magnetic and ferroelectric ordering in hexagonal YMnO3 crystals.[1–5] The coexistence of two or more order parameters in rare-earth manganites provides a wide spectrum of physical responses to electrical, magnetic, and elastic stimuli, making these materials promising candidates for advanced multifunctional electronic devices.[6,7] Rare-earth manganites have two kinds of structural symmetries, which are related to the radius of rare-earth ions (Re3+).[8] ReMnO3 compounds with a larger Re radius (Re = La, Pr, Nd, Dr. D. J. Kim, Dr. T. R. Paudel, Dr. H. Lu, Prof. J. D. Burton, Prof. E. Y. Tsymbal, Prof. A. Gruverman Department of Physics and Astronomy & Nebraska Center for Materials and Nanoscience University of Nebraska Lincoln, NE 68588, USA E-mail: [email protected]; [email protected] J. G. Connell, Prof. S. S. A. Seo Department of Physics and Astronomy University of Kentucky KY 40506, USA
DOI: 10.1002/adma.201403301
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Sm, Eu, Gd, Tb, and Dy) have an orthorhombic structure, while manganites with a smaller Re radius (Re = Ho, Er, Tm, Yb, Lu, and Y) have a hexagonal structure (Figure 1a). The latter group exhibits ferroelectric behavior at temperatures up to 600−1000 K induced by long-range Coulomb interactions and titling of the MnO5 polyhedra.[2] Orthorhombic ReMnO3 compounds with compositions close to the phase boundary, such as TbMnO3, also exhibit ferroelectricity but at much lower temperature (below 30 K).[9] This is due to a change in magnetic order from sinusoidal anti-ferromagnetic to a spin spiral structure occurring at the transition temperature, which breaks the inversion symmetry and brings about a spontaneous polarization along the c-axis. Recently, Lee et al. reported that the hexagonal phase of TbMnO3 could be stabilized through epitaxial strain-engineering by choosing an appropriate substrate.[10] They found that the hexagonal TbMnO3 (h-TbMnO3) films exhibit a much larger polarization (ca. 1 µC/cm2) than the orthorhombic films (0.03−0.08 µC/cm2) and a higher Curie temperature (TC), yet well below room temperature (ca. 60 K). It was suggested that above 60 K the h-TbMnO3 films undergo an anti-ferroelectric phase transition. This conclusion, however, was made based on the dielectric and electrical measurements, which, due to leakage currents, could not be carried out at higher temperatures. In this paper, we perform experimental and theoretical studies of the ferroelectric behavior of strain-engineered TbMnO3 thin films with hexagonal structure. Using local probe characterization techniques that allow electrical measurements at elevated temperatures we demonstrate robust room-temperature ferroelectricity in h-TbMnO3 films, which is accompanied by resistive switching driven by electrically induced polarization reversal. The experimental results are supported by our firstprinciples calculations predicting the ferroelectric polarization in h-TbMnO3 two orders of magnitude larger than that reported in orthorhombic TbMnO3 (o-TbMnO3) at low temperatures. This finding opens up an exciting possibility to strain-engineer a new phase of the rare-earth manganites with novel functionality not available in the bulk, namely room-temperature ferroelectricity and resistive switching. TbMnO3 thin films of ca. 23 nm in thickness were grown epitaxially on Pt(111)/Al2O3(0001) substrates, as described in Materials and Methods. Figure 1b shows a typical X-ray diffraction pattern revealing high film crystallinity. Only (000l) peaks are seen, indicating the hexagonal phase of the TbMnO3 film with the c-axis normal to the substrate. Hexagonal in-plane symmetry, coherent with the Al2O3 substrate, is also confirmed by a phi-scan of X-ray diffraction (Figure S1, Supporting Information). A topographic image shown in Figure 2a reveals
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a nanosized grain structure of the film with a root-mean-square roughness of ca. 1.8 nm. Ferroelectric properties of the h-TbMnO3 films were studied by piezoresponse force microscopy (PFM). Figure 2b shows a local PFM response of the film exhibiting a hysteretic behavior typical for ferroelectric switching. Importantly, hysteretic PFM loops were also observed at an elevated temperature of 180 °C (Figure S2, Supporting Information). To investigate the switching behavior on a larger scale, a 2 µm × 2 µm area of the film was poled by scanning, first, with the negatively, and then, positively biased probe tip. Figure 2c and 2d show the PFM images of the resulting bi-domain pattern, which do not reveal any relaxation for at least 24 hours after poling. We interpret the observation of electrically switchable bistable states in the h-TbMnO3 films as evidence of their ferroelectric state. Although the hysteretic PFM loops and electrically induced reversal in the PFM contrast can occur from electrostatic interactions between the tip and surface charges,[11] we alleviate this effect by performing mechanical switching of polarization via the flexoelectric effect.[12] In this approach, the electrically pre-poled area of the film is scanned with a grounded tip at an incrementally increasing loading force. The stress gradient developed across the film thickness due to the sharp tip apex pressing against the film surface generates the flexoelectric field oriented downward (toward the film-electrode interface). At a certain level of the loading
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Figure 1. Epitaxially stabilized hexagonal TbMnO3 thin films. a) Crystal structure in the ground state of the rare-earth manganites as a function of the radius of the rare-earth elements. b) X-ray diffraction pattern of the h-TbMnO3 film grown on Pt(111)/Al2O3(0001) substrate.
force, the flexoelectric field exceeds the coercive field, which results in polarization reversal from the upward to the downward direction.[12] Figure 2e and 2f show the PFM images of the domain pattern modified by scanning a 1 µm × 1 µm area with the loading force increasing from 150 nN to 1500 nN (corresponding to the applied stress changing from 0.5 to 5 GPa). As seen in Figure 2f, the tip-induced stress leads to a reversal of the PFM phase signal in the area initially poled upward, which is an indication of its polarization reversal to the downward polarization state. Given that no voltage is applied during flexoelectric switching, this result rules out charge deposition or injection at the surface as the origin of the reversal of PFM contrast. A possibility of triboelectric surface charging has been ruled out as well since the mechanically induced switching has been observed while using conductive and non-conductive tips both in the open- and close-circuit conditions as well as, importantly, in sliding and local indentation regimes. Electrically driven ion[13] and oxygen-vacancy[14] migration may also mimic ferroelectric switching behavior. Unfortunately, the flexoelectric effect does not exclude the possibility of oxygen vacancy migration, which could also result in the switchable electromechanical response.[14] Note, however, that strong retention of both polarization states is not consistent with the ion migration mechanism, since oxygen vacancies typically show a preference for accumulation at one of the film interfaces. Electrically induced surface electrochemical reactions as a possible mechanism can also be ruled out on several grounds. For example, in contrast to other reports where electrochemical effect was established,[15,16] a much lower bias has been used in our studies and no topography change after field application has been observed. More importantly, considering that h-ReMnO3 with a Re radius comparable to that of Tb, such as h-HoMnO3 and h-YMnO3, are room-temperature ferroelectrics (see Figure 1a and Figure S3, Supporting Information), it is reasonable to expect that the epitaxially stabilized hexagonal phase of TbMnO3 also exhibits ferroelectric ordering at room temperature. Results of our theoretical modeling and the observed strong electroresistance response, as discussed below, serve as a further evidence of robust ferroelectricity of the h-TbMnO3 thin films at room temperature. Our first-principles calculations reveal that the ferroelectric non-centrosymmetric P63cm phase of h-TbMnO3 is energetically more favorable than the paraelectric centrosymmetric P63/ mmc phase. For both strain-free (fully relaxed) and strained phases (considering strain induced by the Al2O3 substrate), we find an energy difference between paraelectric and ferroelectric phases of ΔE > 20 meV/atom (Table 1), which corresponds to ca. 1160 K/formula unit. This energy gain for the ferroelectric state is very similar to other h-ReMnO3,[3,17] in particular h-YMnO3, where the calculated ΔE of the ferroelectric phase is 21 meV/atom (ca. 1240 K/formula unit) and the experimental TC is 1270 K. Given that TC scales with ΔE, it can be expected that the TC of h-TbMnO3 is similar to that of h-YMnO3,[3] although a Monte Carlo simulation would be required to predict TC explicitly.[18] Berry phase calculations of the ferroelectric phase reveal that the spontaneous polarization is ca. 8 µC/cm2 in the strain-free phase and ca. 18 µC/cm2 in the strained phase, which is two orders of magnitude larger than the experimentally measured polarization of o-TbMnO3 crystals at low
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Figure 2. Testing of ferroelectric switching using PFM. a) Topography and b) local PFM hysteresis loops of the h-TbMnO3 film. c,d) PFM amplitude (c) and phase (d) images of the poled h-TbMnO3 film. The left and right sides were poled by −5 V and +5 V tip biases, respectively. e,f) PFM amplitude (e) and phase (f) images of the same area of the film where the central region was subjected to mechanical stress induced by the probing tip. The tip loading force increases from 150 nN to 1500 nN as indicated by the arrow in (e).
temperature[17] and comparable to the reported values of other h-ReMnO3 single crystals, such as h-YMnO3 (ca. 5.5 µC/cm2) and h-HoMnO3 (ca. 5.6 µC/cm2).[5,19] The polarization of h-TbMnO3 originates from the rotation and tilt of the MnO5 polyhedra similar to those found for other h-ReMnO3 compounds,[2,3] but different from o-TbMnO3 where ferroelectricity is induced by spin frustration.[17] In h-TbMnO3, Tb and Mn ions lie in separate layers as shown in Figure 3a. In the ferroelectric phase, nearby Tb ions move away from their centrosymmetric positions and oxygen ions move in such a way so that the polyhedra, centered at the nearly centrosymmetric Mn ions, tilt away from each other. In this kind of improper ferroelectricity induced by polyhedral rotation and tilt, both cations and anions mostly retain their ionicity. The calculated average Born effective charges Z* for Tb, Mn and O, defined as the displacement derivative of polarization at zero electric field, are 3.9, 3.7, and −2.5, respectively, which are close to their formal oxidation states of +3, +3, and −2, respectively. This is due to the absence of charge transfer between cations, as seen from spin-resolved density of states (Figure 3b). In contrast, the Table 1. Lattice parameters, energy and polarization of hexagonal TbMnO3 films. Lattice constant of c-axis [Å]
ΔE [meV/atom]
6.20 (strain-free)
11.53
−24
7.8
6.27 (ref.[10])
11.46
−24
7.9
6.76 (strained)
10.57
−22
18
Lattice constant of a-axis [Å]
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PS [µC/cm2]
Z* values in displacive ferroelectrics (e.g., BaTiO3) are more than a factor of two larger than the formal ionic charges. The displacements leading to the ferroelectric phase consist of four phonon modes of the undistorted paraelectric phase: G1+, G2−, K3, and K1, as explained in the Supporting Information and shown in Figure S4. The large lattice instability of the paraelectric phase is driven by the K3 mode, which displays a well-pronounced double-well behavior of the phonon mode energy (Figure 3c, bottom panel). On the contrary, the polar G2− and non-polar K1 modes show single minima around the paraelectric state (Figure 3c, middle and top panels, respectively). The G1+ mode involves breathing displacements and thus cannot lead to symmetry lowering. In spite of the double-well potential, however, it is not the K3 mode alone that is responsible for spontaneous polarization. Similar to h-YMnO3,[3] we find that polarization of h-TbMnO3 is controlled by the interplay between the G2− and K3 modes. As evident from Figure 3d, in the presence of the K3 mode, the center of symmetry of the G2− mode shifts progressively from the origin, reducing its energy with increasing contribution from the K3 mode. In the ground state, where all the phonon modes are contributing (see Table S1, Supporting Information), the net spontaneous polarization of 8 µC/cm2 (in the strain-free state) is largely determined by the G2− mode, whereas the energy gain is mainly due to the K3 mode. Next, we investigate the resistive switching behavior of the h-TbMnO3 films, expecting that resistance through the film should be dependent on polarization.[5] Figure 4a shows a local conductance map of the h-TbMnO3 film obtained with a −1.5 V read bias (which is well below the coercive voltage of the film
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COMMUNICATION Figure 3. Theoretical modeling. a) Spin-resolved density of states (DOS) of O, Tb, and Mn ions in the ferroelectric phase compared to those in the paraelectric phase. The black and red lines represent spin-up and spin-down DOS of the ferroelectric phase, respectively, and the filled curves behind represent the spin-resolved DOS of the paraelectric phase. b) Calculated ferroelectric P63cm crystal structure of bulk h-TbMnO3. The arrows indicate the direction of shift of ions during paraelectric to ferroelectric phase transition. In the ferroelectric phase, the mirror plane perpendicular to the c-axis is lost so that the O polyhedra around the Mn ions in the triangular lattice tilt away from each other. c) Energy of the K1, Γ2−, and K3 phonon modes (top, middle, and bottom panels respectively) as a function of mode amplitude, indicating the lattice instability driven by the K3 mode. d) Energy of the polar Γ2− mode in with the indicated amplitude of the K3 mode frozen in. The energy minimum of the Γ2− mode shifts from the origin in the presence of the K3 mode, thereby reducing the cross cancellation of the Tb-centered dipoles and thus leading to an increase of polarization.
in Figure 2b) after the upper and lower areas were poled by the positively and negatively biased tip, respectively. The intermediate contrast of the non-poled region could be associated with a polydomain structure (not discernible in PFM due to the small domain size) in the as-grown film that would yield a conductance level different from the negatively and positively poled domains. Figure 4b shows representative current–voltage (I–V) curves measured in the low resistance (ON) and the high resistance (OFF) areas. It is seen that the resistance changes by more than two orders of magnitude upon polarization reversal, which is comparable with the resistive switching effects reported for other ferroelectric systems.[20–25] The observed conduction behavior cannot be attributed to quantum tunneling[26] since the film is too thick (ca. 23 nm).
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The I–V characteristics are of semiconducting nature, which is consistent with the fact that h-TbMnO3 has a relatively narrow band gap of ca. 1.4 eV.[27] Interestingly, in spite of the significant spatial variations in local conductance (see Figure S5, Supporting Information), normalized I–V curves of all resistance states can be fit to a single universal curve (Figure 4c), implying a single conduction mechanism. Given the symmetric appearance of the I–V curves, we propose a back-to-back Schottky diodes model to explain the conduction behavior in semiconducting h-TbMnO3 films.[28] In this model, there are two important parameters: the Schottky barrier height ΦB, and an ideality factor n of the interfaces. The ideality factor reflects non-ideal effects such as interface states, image-forces, and generationrecombination in the space charge region. It is equal to 1 for
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Figure 4. Polarization-controlled resistive switching. a) A local current map of the h-TbMnO3 film after poling with +6 V (top) and −6 V (bottom) tip bias and as-grown (unpoled) film (middle). The image was acquired with a DC read voltage of −1.5 V. The arrows indicate the locations where the crosssection line profiles in Figure S5 in the Supporting Information were taken. b) (Top panel) Representative local current–voltage (I–V) curves measured in the low resistance (ON) area poled by +6 V and the high resistance (OFF) area poled by −6 V; (bottom panel) OFF/ON resistance ratio curve calculated from the I–V curves in (b). c) Normalized I–V curves measured in positively and negatively poled areas and as-grown (unpoled) TbMnO3 film. The red dashed line is a fitting result using the back-to-back Schottky diodes model with the ideality factor of n = 1.09 and an arbitrary barrier height ΦB.
an ideal Schottky contact, but is larger than 1 in non-ideal cases (normally, 1 < n < 1.2). The back-to-back Schottky diode model well describes the I–V curves for all resistance states (the red dashed curve in Figure 4c) and the fitting yields the same value of n for both interfaces. This is reasonable since we used a Pt-coated tip for the film on Pt bottom electrode. Although the exact value of the Schottky barrier height could not be determined, we modeled the I–V curves for certain values of ΦB and n. In the case of fixed n and variable ΦB, all normalized I–V curves can be fit by a single curve, while varying n at a fixed value of ΦB yields different I–V curves (see Figure S6, Supporting Information). This analysis suggests that the conduction of every resistance state in the h-TbMnO3 films is determined by the Schottky barrier height, which depends on polarization direction. The observed resistive switching driven by ferroelectric polarization reversal further confirms the existence of ferroelectricity in h-TbMnO3 films at room temperature. This finding is at odds with the earlier reported ferroelectric–anti-ferroelectric transition in h-TbMnO3 films occurring at ca. 60 K.[10] We note, however, that this result was obtained by temperaturedependent measurements of polarization hysteresis loops. The anti-ferroelectric-like behavior of h-TbMnO3 in ref.[10] can be explained by defect dipole alignments during the film deposition at high temperature, which are not easy to reverse at room or lower temperatures.[29,30] It is possible that the h-TbMnO3 films in ref.[10] were not anti-ferroelectric but rather ferroelectric exhibiting spontaneous back-switching of polarization due to aligned defect dipoles. In addition, the polarization measurements could be affected by increased leakage current due to point defects or vacancies of any elemental components of the film material. In the PFM experimental approach, used in our studies, a small contact tip-sample area allows application of a voltage pulse of a much larger amplitude without causing much leakage current, thereby allowing detection of the ferroelectric behavior of h-TbMnO3 films at temperatures well above the room temperature.
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Our experimental and theoretical results provide strong evidence for room-temperature ferroelectricity in epitaxially stabilized hexagonal TbMnO3 thin films, demonstrating that we could strain-engineer new phases and properties of functional materials using epitaxial growth. Given that hexagonal TbMnO3 is a narrow band gap semiconductor, ferroelectricity of this material at room temperature allows polarization-dependent resistive switching with thicker barriers compared to those required for ferroelectric tunnel junctions, thus, making the fabrication process much easier. This property, in conjunction with the reported photovoltaic behavior,[27] makes hexagonal TbMnO3 films promising materials for novel multifunctional electronic devices.
Experimental Section Sample Preparation: Epitaxial h-TbMnO3 thin films with the thickness of 20 unit cells (ca. 23 nm) were grown on Pt(111)/Al2O3(0001) substrates using pulsed laser deposition at a substrate temperature of 800 °C, oxygen partial pressure of 30 mTorr, and a KrF excimer (λ = 248 nm) laser fluence of 1.2 J/cm2.[31] The Pt bottom layers were deposited on Al2O3 (0001) single crystal substrates at 600 °C using rf-sputtering. X-ray diffraction scans in Figure 1b confirmed hexagonal structure of the epitaxially grown films. Data on Figure S1 in the Supporting Information illustrate the in-plane epitaxy and six-fold symmetry of the (0001)-oriented h-TbMnO3 thin films. There was no evidence of any other phases, such as TbMn2O5, in the X-ray diffraction. Electrical Characterization: An atomic force microscope (Asylum MFP-3D) operating in the resonance-enhanced PFM and conductive atomic force microscopy modes was used to test the ferroelectric and conducting properties of the films, respectively. Conductive Pt-coated cantilevers (Nanosensors PPP-EFM) were used for PFM measurements. Mechanical switching of the polarization in the h-TbMnO3 film was performed by scanning with a grounded tip at a 150–1500 nN loading force.[12] Local piezoelectric hysteresis loops were obtained at several fixed locations by the pulsed switching scheme. Local conductance maps and current–voltage curves were obtained using electrically biased Pt-coated tips. Theoretical Modeling: Theoretical modeling of the h-TbMnO3 films was performed using density functional theory, the projected augmented wave
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Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
Acknowledgements This research was primarily supported by the US Department of Energy, Office of Science, Basic Energy Sciences, Materials Sciences and Engineering Division, under Award DE-SC0004876 (electrical characterization and theoretical modeling). Computations were performed utilizing the Holland Computing Center of the University of Nebraska and the Center for Nanophase Materials Sciences, which is sponsored at Oak Ridge National Laboratory by the Scientific User Facilities Division, Office of Basic Energy Sciences, US Department of Energy. Structural characterization at the University of NebraskaLincoln was supported by the National Science Foundation through Materials Research Science and Engineering Center under Grant DMR0820521. The work at the University of Kentucky (sample preparation) was supported by the National Science Foundation through Grant EPS-0814194 (the Center for Advanced Materials) and by the Kentucky Science and Engineering Foundation with the Kentucky Science and Technology Corporation through Grant Agreement KSEF-148–502–12– 303. The authors acknowledge discussion with Dr. Jeevaka Weerasinghe regarding the prediction of Tc using a Monte Carlo simulation and with Dr. Sitaram Jaswal regarding the group-theory-based phonon-mode analysis. The authors thank Prof. Peter Dowben, Iori Tanabe, Timothy Vo, and Prof. Alexander Sinitskii for the Raman spectroscopy measurements. Received: July 22, 2014 Revised: September 2, 2014 Published online: October 18, 2014
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method, and the Perdew–Burke–Ernzerhof functional, as implemented in Vienna ab initio simulation package.[32] We fully relaxed the structure with the force convergence limit of 0.005 eV/atom. Correlation effects beyond generalized gradient approximation (GGA) were treated at a semiempirical GGA+U level within a rotationally invariant formalism with U = 1.5 eV for the Mn 3d-orbitals. Neighboring spin moments of ca. 3.72 µB/Mn were aligned anti-ferromagnetically both in-plane and out-ofplane, resulting in zero net magnetic moment in collinear approximation. The calculated lattice constants of the strain-free h-TbMnO3 are very similar to those found by X-ray diffraction measurements.[10]
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Room-Temperature Ferroelectricity in Hexagonal TbMnO3 Thin Films Dong Jik Kim,* Tula R. Paudel, Haidong Lu, John D. Burton, John G. Connell, Evgeny Y. Tsymbal, S. S. Ambrose Seo, and Alexei Gruverman* Rare-earth manganite materials (ReMnO3) exhibit a wealth of fundamentally exciting and technologically appealing properties. For example, a coupled magnetic and ferroelectric ordering in these materials allows the manipulation of ferroelectric polarization with a magnetic field and/or magnetization with an electric field, thereby opening new possibilities for development of multifunctional electronic devices. A fundamentally intriguing aspect is that, depending on the size of the Re ion, these materials exhibit either an orthorhombic or a hexagonal structure. While the orthorhombic rare-earth manganites, such as TbMnO3, demonstrate ferroelectric ordering at rather low temperatures (well below 100 K), the hexagonal ReMnO3 tend to be high-temperature ferroelectrics with a transition temperature typically above 600 K. Here, we report room-temperature ferroelectricity in thin films of epitaxially stabilized hexagonal TbMnO3, which in the bulk form exhibits orthorhombic structure. Our theoretical modeling reveals that the ferroelectric polarization of hexagonal TbMnO3 arises from the rotation and tilt of the MnO5 polyhedra, similar to other hexagonal manganites. Switchable ferroelectricity is accompanied by significant polarization-dependent resistive switching, which is attributed to polarization-dependent Schottky barrier modulation. The obtained results demonstrate that strain-engineering of rareearth manganites is a viable route to realize new physical properties of these materials and achieve novel functionalities at elevated temperatures useful for practical applications. Rare-earth manganites (ReMnO3) have attracted much attention since the discovery of coupled magnetic and ferroelectric ordering in hexagonal YMnO3 crystals.[1–5] The coexistence of two or more order parameters in rare-earth manganites provides a wide spectrum of physical responses to electrical, magnetic, and elastic stimuli, making these materials promising candidates for advanced multifunctional electronic devices.[6,7] Rare-earth manganites have two kinds of structural symmetries, which are related to the radius of rare-earth ions (Re3+).[8] ReMnO3 compounds with a larger Re radius (Re = La, Pr, Nd, Dr. D. J. Kim, Dr. T. R. Paudel, Dr. H. Lu, Prof. J. D. Burton, Prof. E. Y. Tsymbal, Prof. A. Gruverman Department of Physics and Astronomy & Nebraska Center for Materials and Nanoscience University of Nebraska Lincoln, NE 68588, USA E-mail: [email protected]; [email protected] J. G. Connell, Prof. S. S. A. Seo Department of Physics and Astronomy University of Kentucky KY 40506, USA
DOI: 10.1002/adma.201403301
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Sm, Eu, Gd, Tb, and Dy) have an orthorhombic structure, while manganites with a smaller Re radius (Re = Ho, Er, Tm, Yb, Lu, and Y) have a hexagonal structure (Figure 1a). The latter group exhibits ferroelectric behavior at temperatures up to 600−1000 K induced by long-range Coulomb interactions and titling of the MnO5 polyhedra.[2] Orthorhombic ReMnO3 compounds with compositions close to the phase boundary, such as TbMnO3, also exhibit ferroelectricity but at much lower temperature (below 30 K).[9] This is due to a change in magnetic order from sinusoidal anti-ferromagnetic to a spin spiral structure occurring at the transition temperature, which breaks the inversion symmetry and brings about a spontaneous polarization along the c-axis. Recently, Lee et al. reported that the hexagonal phase of TbMnO3 could be stabilized through epitaxial strain-engineering by choosing an appropriate substrate.[10] They found that the hexagonal TbMnO3 (h-TbMnO3) films exhibit a much larger polarization (ca. 1 µC/cm2) than the orthorhombic films (0.03−0.08 µC/cm2) and a higher Curie temperature (TC), yet well below room temperature (ca. 60 K). It was suggested that above 60 K the h-TbMnO3 films undergo an anti-ferroelectric phase transition. This conclusion, however, was made based on the dielectric and electrical measurements, which, due to leakage currents, could not be carried out at higher temperatures. In this paper, we perform experimental and theoretical studies of the ferroelectric behavior of strain-engineered TbMnO3 thin films with hexagonal structure. Using local probe characterization techniques that allow electrical measurements at elevated temperatures we demonstrate robust room-temperature ferroelectricity in h-TbMnO3 films, which is accompanied by resistive switching driven by electrically induced polarization reversal. The experimental results are supported by our firstprinciples calculations predicting the ferroelectric polarization in h-TbMnO3 two orders of magnitude larger than that reported in orthorhombic TbMnO3 (o-TbMnO3) at low temperatures. This finding opens up an exciting possibility to strain-engineer a new phase of the rare-earth manganites with novel functionality not available in the bulk, namely room-temperature ferroelectricity and resistive switching. TbMnO3 thin films of ca. 23 nm in thickness were grown epitaxially on Pt(111)/Al2O3(0001) substrates, as described in Materials and Methods. Figure 1b shows a typical X-ray diffraction pattern revealing high film crystallinity. Only (000l) peaks are seen, indicating the hexagonal phase of the TbMnO3 film with the c-axis normal to the substrate. Hexagonal in-plane symmetry, coherent with the Al2O3 substrate, is also confirmed by a phi-scan of X-ray diffraction (Figure S1, Supporting Information). A topographic image shown in Figure 2a reveals
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a nanosized grain structure of the film with a root-mean-square roughness of ca. 1.8 nm. Ferroelectric properties of the h-TbMnO3 films were studied by piezoresponse force microscopy (PFM). Figure 2b shows a local PFM response of the film exhibiting a hysteretic behavior typical for ferroelectric switching. Importantly, hysteretic PFM loops were also observed at an elevated temperature of 180 °C (Figure S2, Supporting Information). To investigate the switching behavior on a larger scale, a 2 µm × 2 µm area of the film was poled by scanning, first, with the negatively, and then, positively biased probe tip. Figure 2c and 2d show the PFM images of the resulting bi-domain pattern, which do not reveal any relaxation for at least 24 hours after poling. We interpret the observation of electrically switchable bistable states in the h-TbMnO3 films as evidence of their ferroelectric state. Although the hysteretic PFM loops and electrically induced reversal in the PFM contrast can occur from electrostatic interactions between the tip and surface charges,[11] we alleviate this effect by performing mechanical switching of polarization via the flexoelectric effect.[12] In this approach, the electrically pre-poled area of the film is scanned with a grounded tip at an incrementally increasing loading force. The stress gradient developed across the film thickness due to the sharp tip apex pressing against the film surface generates the flexoelectric field oriented downward (toward the film-electrode interface). At a certain level of the loading
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Figure 1. Epitaxially stabilized hexagonal TbMnO3 thin films. a) Crystal structure in the ground state of the rare-earth manganites as a function of the radius of the rare-earth elements. b) X-ray diffraction pattern of the h-TbMnO3 film grown on Pt(111)/Al2O3(0001) substrate.
force, the flexoelectric field exceeds the coercive field, which results in polarization reversal from the upward to the downward direction.[12] Figure 2e and 2f show the PFM images of the domain pattern modified by scanning a 1 µm × 1 µm area with the loading force increasing from 150 nN to 1500 nN (corresponding to the applied stress changing from 0.5 to 5 GPa). As seen in Figure 2f, the tip-induced stress leads to a reversal of the PFM phase signal in the area initially poled upward, which is an indication of its polarization reversal to the downward polarization state. Given that no voltage is applied during flexoelectric switching, this result rules out charge deposition or injection at the surface as the origin of the reversal of PFM contrast. A possibility of triboelectric surface charging has been ruled out as well since the mechanically induced switching has been observed while using conductive and non-conductive tips both in the open- and close-circuit conditions as well as, importantly, in sliding and local indentation regimes. Electrically driven ion[13] and oxygen-vacancy[14] migration may also mimic ferroelectric switching behavior. Unfortunately, the flexoelectric effect does not exclude the possibility of oxygen vacancy migration, which could also result in the switchable electromechanical response.[14] Note, however, that strong retention of both polarization states is not consistent with the ion migration mechanism, since oxygen vacancies typically show a preference for accumulation at one of the film interfaces. Electrically induced surface electrochemical reactions as a possible mechanism can also be ruled out on several grounds. For example, in contrast to other reports where electrochemical effect was established,[15,16] a much lower bias has been used in our studies and no topography change after field application has been observed. More importantly, considering that h-ReMnO3 with a Re radius comparable to that of Tb, such as h-HoMnO3 and h-YMnO3, are room-temperature ferroelectrics (see Figure 1a and Figure S3, Supporting Information), it is reasonable to expect that the epitaxially stabilized hexagonal phase of TbMnO3 also exhibits ferroelectric ordering at room temperature. Results of our theoretical modeling and the observed strong electroresistance response, as discussed below, serve as a further evidence of robust ferroelectricity of the h-TbMnO3 thin films at room temperature. Our first-principles calculations reveal that the ferroelectric non-centrosymmetric P63cm phase of h-TbMnO3 is energetically more favorable than the paraelectric centrosymmetric P63/ mmc phase. For both strain-free (fully relaxed) and strained phases (considering strain induced by the Al2O3 substrate), we find an energy difference between paraelectric and ferroelectric phases of ΔE > 20 meV/atom (Table 1), which corresponds to ca. 1160 K/formula unit. This energy gain for the ferroelectric state is very similar to other h-ReMnO3,[3,17] in particular h-YMnO3, where the calculated ΔE of the ferroelectric phase is 21 meV/atom (ca. 1240 K/formula unit) and the experimental TC is 1270 K. Given that TC scales with ΔE, it can be expected that the TC of h-TbMnO3 is similar to that of h-YMnO3,[3] although a Monte Carlo simulation would be required to predict TC explicitly.[18] Berry phase calculations of the ferroelectric phase reveal that the spontaneous polarization is ca. 8 µC/cm2 in the strain-free phase and ca. 18 µC/cm2 in the strained phase, which is two orders of magnitude larger than the experimentally measured polarization of o-TbMnO3 crystals at low
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Figure 2. Testing of ferroelectric switching using PFM. a) Topography and b) local PFM hysteresis loops of the h-TbMnO3 film. c,d) PFM amplitude (c) and phase (d) images of the poled h-TbMnO3 film. The left and right sides were poled by −5 V and +5 V tip biases, respectively. e,f) PFM amplitude (e) and phase (f) images of the same area of the film where the central region was subjected to mechanical stress induced by the probing tip. The tip loading force increases from 150 nN to 1500 nN as indicated by the arrow in (e).
temperature[17] and comparable to the reported values of other h-ReMnO3 single crystals, such as h-YMnO3 (ca. 5.5 µC/cm2) and h-HoMnO3 (ca. 5.6 µC/cm2).[5,19] The polarization of h-TbMnO3 originates from the rotation and tilt of the MnO5 polyhedra similar to those found for other h-ReMnO3 compounds,[2,3] but different from o-TbMnO3 where ferroelectricity is induced by spin frustration.[17] In h-TbMnO3, Tb and Mn ions lie in separate layers as shown in Figure 3a. In the ferroelectric phase, nearby Tb ions move away from their centrosymmetric positions and oxygen ions move in such a way so that the polyhedra, centered at the nearly centrosymmetric Mn ions, tilt away from each other. In this kind of improper ferroelectricity induced by polyhedral rotation and tilt, both cations and anions mostly retain their ionicity. The calculated average Born effective charges Z* for Tb, Mn and O, defined as the displacement derivative of polarization at zero electric field, are 3.9, 3.7, and −2.5, respectively, which are close to their formal oxidation states of +3, +3, and −2, respectively. This is due to the absence of charge transfer between cations, as seen from spin-resolved density of states (Figure 3b). In contrast, the Table 1. Lattice parameters, energy and polarization of hexagonal TbMnO3 films. Lattice constant of c-axis [Å]
ΔE [meV/atom]
6.20 (strain-free)
11.53
−24
7.8
6.27 (ref.[10])
11.46
−24
7.9
6.76 (strained)
10.57
−22
18
Lattice constant of a-axis [Å]
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PS [µC/cm2]
Z* values in displacive ferroelectrics (e.g., BaTiO3) are more than a factor of two larger than the formal ionic charges. The displacements leading to the ferroelectric phase consist of four phonon modes of the undistorted paraelectric phase: G1+, G2−, K3, and K1, as explained in the Supporting Information and shown in Figure S4. The large lattice instability of the paraelectric phase is driven by the K3 mode, which displays a well-pronounced double-well behavior of the phonon mode energy (Figure 3c, bottom panel). On the contrary, the polar G2− and non-polar K1 modes show single minima around the paraelectric state (Figure 3c, middle and top panels, respectively). The G1+ mode involves breathing displacements and thus cannot lead to symmetry lowering. In spite of the double-well potential, however, it is not the K3 mode alone that is responsible for spontaneous polarization. Similar to h-YMnO3,[3] we find that polarization of h-TbMnO3 is controlled by the interplay between the G2− and K3 modes. As evident from Figure 3d, in the presence of the K3 mode, the center of symmetry of the G2− mode shifts progressively from the origin, reducing its energy with increasing contribution from the K3 mode. In the ground state, where all the phonon modes are contributing (see Table S1, Supporting Information), the net spontaneous polarization of 8 µC/cm2 (in the strain-free state) is largely determined by the G2− mode, whereas the energy gain is mainly due to the K3 mode. Next, we investigate the resistive switching behavior of the h-TbMnO3 films, expecting that resistance through the film should be dependent on polarization.[5] Figure 4a shows a local conductance map of the h-TbMnO3 film obtained with a −1.5 V read bias (which is well below the coercive voltage of the film
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COMMUNICATION Figure 3. Theoretical modeling. a) Spin-resolved density of states (DOS) of O, Tb, and Mn ions in the ferroelectric phase compared to those in the paraelectric phase. The black and red lines represent spin-up and spin-down DOS of the ferroelectric phase, respectively, and the filled curves behind represent the spin-resolved DOS of the paraelectric phase. b) Calculated ferroelectric P63cm crystal structure of bulk h-TbMnO3. The arrows indicate the direction of shift of ions during paraelectric to ferroelectric phase transition. In the ferroelectric phase, the mirror plane perpendicular to the c-axis is lost so that the O polyhedra around the Mn ions in the triangular lattice tilt away from each other. c) Energy of the K1, Γ2−, and K3 phonon modes (top, middle, and bottom panels respectively) as a function of mode amplitude, indicating the lattice instability driven by the K3 mode. d) Energy of the polar Γ2− mode in with the indicated amplitude of the K3 mode frozen in. The energy minimum of the Γ2− mode shifts from the origin in the presence of the K3 mode, thereby reducing the cross cancellation of the Tb-centered dipoles and thus leading to an increase of polarization.
in Figure 2b) after the upper and lower areas were poled by the positively and negatively biased tip, respectively. The intermediate contrast of the non-poled region could be associated with a polydomain structure (not discernible in PFM due to the small domain size) in the as-grown film that would yield a conductance level different from the negatively and positively poled domains. Figure 4b shows representative current–voltage (I–V) curves measured in the low resistance (ON) and the high resistance (OFF) areas. It is seen that the resistance changes by more than two orders of magnitude upon polarization reversal, which is comparable with the resistive switching effects reported for other ferroelectric systems.[20–25] The observed conduction behavior cannot be attributed to quantum tunneling[26] since the film is too thick (ca. 23 nm).
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The I–V characteristics are of semiconducting nature, which is consistent with the fact that h-TbMnO3 has a relatively narrow band gap of ca. 1.4 eV.[27] Interestingly, in spite of the significant spatial variations in local conductance (see Figure S5, Supporting Information), normalized I–V curves of all resistance states can be fit to a single universal curve (Figure 4c), implying a single conduction mechanism. Given the symmetric appearance of the I–V curves, we propose a back-to-back Schottky diodes model to explain the conduction behavior in semiconducting h-TbMnO3 films.[28] In this model, there are two important parameters: the Schottky barrier height ΦB, and an ideality factor n of the interfaces. The ideality factor reflects non-ideal effects such as interface states, image-forces, and generationrecombination in the space charge region. It is equal to 1 for
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Figure 4. Polarization-controlled resistive switching. a) A local current map of the h-TbMnO3 film after poling with +6 V (top) and −6 V (bottom) tip bias and as-grown (unpoled) film (middle). The image was acquired with a DC read voltage of −1.5 V. The arrows indicate the locations where the crosssection line profiles in Figure S5 in the Supporting Information were taken. b) (Top panel) Representative local current–voltage (I–V) curves measured in the low resistance (ON) area poled by +6 V and the high resistance (OFF) area poled by −6 V; (bottom panel) OFF/ON resistance ratio curve calculated from the I–V curves in (b). c) Normalized I–V curves measured in positively and negatively poled areas and as-grown (unpoled) TbMnO3 film. The red dashed line is a fitting result using the back-to-back Schottky diodes model with the ideality factor of n = 1.09 and an arbitrary barrier height ΦB.
an ideal Schottky contact, but is larger than 1 in non-ideal cases (normally, 1 < n < 1.2). The back-to-back Schottky diode model well describes the I–V curves for all resistance states (the red dashed curve in Figure 4c) and the fitting yields the same value of n for both interfaces. This is reasonable since we used a Pt-coated tip for the film on Pt bottom electrode. Although the exact value of the Schottky barrier height could not be determined, we modeled the I–V curves for certain values of ΦB and n. In the case of fixed n and variable ΦB, all normalized I–V curves can be fit by a single curve, while varying n at a fixed value of ΦB yields different I–V curves (see Figure S6, Supporting Information). This analysis suggests that the conduction of every resistance state in the h-TbMnO3 films is determined by the Schottky barrier height, which depends on polarization direction. The observed resistive switching driven by ferroelectric polarization reversal further confirms the existence of ferroelectricity in h-TbMnO3 films at room temperature. This finding is at odds with the earlier reported ferroelectric–anti-ferroelectric transition in h-TbMnO3 films occurring at ca. 60 K.[10] We note, however, that this result was obtained by temperaturedependent measurements of polarization hysteresis loops. The anti-ferroelectric-like behavior of h-TbMnO3 in ref.[10] can be explained by defect dipole alignments during the film deposition at high temperature, which are not easy to reverse at room or lower temperatures.[29,30] It is possible that the h-TbMnO3 films in ref.[10] were not anti-ferroelectric but rather ferroelectric exhibiting spontaneous back-switching of polarization due to aligned defect dipoles. In addition, the polarization measurements could be affected by increased leakage current due to point defects or vacancies of any elemental components of the film material. In the PFM experimental approach, used in our studies, a small contact tip-sample area allows application of a voltage pulse of a much larger amplitude without causing much leakage current, thereby allowing detection of the ferroelectric behavior of h-TbMnO3 films at temperatures well above the room temperature.
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Our experimental and theoretical results provide strong evidence for room-temperature ferroelectricity in epitaxially stabilized hexagonal TbMnO3 thin films, demonstrating that we could strain-engineer new phases and properties of functional materials using epitaxial growth. Given that hexagonal TbMnO3 is a narrow band gap semiconductor, ferroelectricity of this material at room temperature allows polarization-dependent resistive switching with thicker barriers compared to those required for ferroelectric tunnel junctions, thus, making the fabrication process much easier. This property, in conjunction with the reported photovoltaic behavior,[27] makes hexagonal TbMnO3 films promising materials for novel multifunctional electronic devices.
Experimental Section Sample Preparation: Epitaxial h-TbMnO3 thin films with the thickness of 20 unit cells (ca. 23 nm) were grown on Pt(111)/Al2O3(0001) substrates using pulsed laser deposition at a substrate temperature of 800 °C, oxygen partial pressure of 30 mTorr, and a KrF excimer (λ = 248 nm) laser fluence of 1.2 J/cm2.[31] The Pt bottom layers were deposited on Al2O3 (0001) single crystal substrates at 600 °C using rf-sputtering. X-ray diffraction scans in Figure 1b confirmed hexagonal structure of the epitaxially grown films. Data on Figure S1 in the Supporting Information illustrate the in-plane epitaxy and six-fold symmetry of the (0001)-oriented h-TbMnO3 thin films. There was no evidence of any other phases, such as TbMn2O5, in the X-ray diffraction. Electrical Characterization: An atomic force microscope (Asylum MFP-3D) operating in the resonance-enhanced PFM and conductive atomic force microscopy modes was used to test the ferroelectric and conducting properties of the films, respectively. Conductive Pt-coated cantilevers (Nanosensors PPP-EFM) were used for PFM measurements. Mechanical switching of the polarization in the h-TbMnO3 film was performed by scanning with a grounded tip at a 150–1500 nN loading force.[12] Local piezoelectric hysteresis loops were obtained at several fixed locations by the pulsed switching scheme. Local conductance maps and current–voltage curves were obtained using electrically biased Pt-coated tips. Theoretical Modeling: Theoretical modeling of the h-TbMnO3 films was performed using density functional theory, the projected augmented wave
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Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
Acknowledgements This research was primarily supported by the US Department of Energy, Office of Science, Basic Energy Sciences, Materials Sciences and Engineering Division, under Award DE-SC0004876 (electrical characterization and theoretical modeling). Computations were performed utilizing the Holland Computing Center of the University of Nebraska and the Center for Nanophase Materials Sciences, which is sponsored at Oak Ridge National Laboratory by the Scientific User Facilities Division, Office of Basic Energy Sciences, US Department of Energy. Structural characterization at the University of NebraskaLincoln was supported by the National Science Foundation through Materials Research Science and Engineering Center under Grant DMR0820521. The work at the University of Kentucky (sample preparation) was supported by the National Science Foundation through Grant EPS-0814194 (the Center for Advanced Materials) and by the Kentucky Science and Engineering Foundation with the Kentucky Science and Technology Corporation through Grant Agreement KSEF-148–502–12– 303. The authors acknowledge discussion with Dr. Jeevaka Weerasinghe regarding the prediction of Tc using a Monte Carlo simulation and with Dr. Sitaram Jaswal regarding the group-theory-based phonon-mode analysis. The authors thank Prof. Peter Dowben, Iori Tanabe, Timothy Vo, and Prof. Alexander Sinitskii for the Raman spectroscopy measurements. Received: July 22, 2014 Revised: September 2, 2014 Published online: October 18, 2014
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method, and the Perdew–Burke–Ernzerhof functional, as implemented in Vienna ab initio simulation package.[32] We fully relaxed the structure with the force convergence limit of 0.005 eV/atom. Correlation effects beyond generalized gradient approximation (GGA) were treated at a semiempirical GGA+U level within a rotationally invariant formalism with U = 1.5 eV for the Mn 3d-orbitals. Neighboring spin moments of ca. 3.72 µB/Mn were aligned anti-ferromagnetically both in-plane and out-ofplane, resulting in zero net magnetic moment in collinear approximation. The calculated lattice constants of the strain-free h-TbMnO3 are very similar to those found by X-ray diffraction measurements.[10]
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