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The essentials for the control of charge-orbital ordering in thin films of perovskite manganites
Materials Science and Engineering B 173 (2010) 51–56
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Materials Science and Engineering B journal homepage: www.elsevier.com/locate/mseb
The essentials for the control of charge-orbital ordering in thin films of perovskite manganites Yasushi Ogimoto a,b,c,∗ , Masao Nakamura d,1 , Nobutaka Harada d , Naoki Ogawa a,b , Kenjiro Miyano a,b a
Research Center for Advanced Science and Technology (RCAST), The University of Tokyo, Tokyo 153-8904, Japan CREST, Japan Science and Technology Agency, Saitama 332-0012, Japan Fuji Electric Advanced Technology Co., Ltd., Tokyo 191-8502, Japan d Department of Applied Physics, The University of Tokyo, Tokyo 113-8656, Japan b c
a r t i c l e
i n f o
Article history: Received 1 July 2009 Received in revised form 1 February 2010 Accepted 3 February 2010 Keywords: Charge-orbital ordering Phase transition Epitaxial strain Correlated electron interface
a b s t r a c t Gigantic responses arising from the melting of charge-orbital ordering (COO) in perovskite manganites have been attractive both in material science and technologies. We review the essentials to realize such phenomena in a thin film form aiming at a possible application to correlated electron devices, namely, the role of epitaxial strain on (1 1 0) substrates. The crucial issues in designing the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0) are also discussed. In addition, the epitaxial growth of high-Tco materials such as (Bi, Sr)MnO3 and A-site ordered SmBaMn2 O6 thin films is described, which is indispensable towards room temperature operation of correlated electron devices. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Gigantic and ultrafast responses arising from the melting of charge-orbital ordering (COO) in perovskite manganites have been attractive both in material science and technologies, because of the rich physics underlying an antiferromagnetic insulator (AFMI) to a ferromagnetic metal (FMM) transition and its possible applicability to correlated electron devices [1,2]. The transition between the FMM and the COO states is first-order, and is coupled to the lattice degree of freedom through the Jahn–Teller (J–T) interaction. Therefore, it is of crucial importance to have a deep understanding of an epitaxial strain from the substrates to control the COO and its AFMI–FMM transition in a thin film form. In the case of the films coherently grown on (0 0 1) perovskite substrates, Konishi et al. demonstrated that the ground states of La0.5 Sr0.5 MnO3 epitaxial thin films are drastically modulated by the substrate-induced strain through the OO close to the phase boundaries adjoining Atype AFMI, FMM, and C-type AFMI states [3]. This striking example also means that the first-order phase transition associated with the COO is prohibited for the epitaxial films coherently grown on (0 0 1)
∗ Corresponding author at: Research Center for Advanced Science and Technology (RCAST), The University of Tokyo, Photonics Materials, 4-6-1, Komaba, Meguro-ku, Tokyo 153-8904, Japan. Tel.: +81 3 5452 5076; fax: +81 3 5452 5076. E-mail addresses: y [email protected], y [email protected] (Y. Ogimoto). 1 Present address: Cross-correlated Materials Research Group (CMRG), RIKEN, Saitama 351-0198, Japan. 0921-5107/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2010.02.009
substrates. The difficulty arises from severe restriction on the lattice deformation in which the in-plane lattice is clamped to the substrate, and the type of which is limited to an orthogonal-mode due to the tetragonal symmetry on (0 0 1) substrates [4]. To escape from the strain effect above, it is easy to make a relaxed film utilizing a large lattice mismatch between the films and the substrates [5]. Since the in-plane lattice is not restricted to the substrates in the relaxed films, the lattice deformation associated with the I–M transition is possible similar to that of the bulk sample, which could lead to the reproduction of the bulk properties [5,6]. However, the films are polycrystalline with a large amount of defects (e.g., misfit dislocations and grain boundaries) which will make the I–M transition diffusive due to the disorder effect [7–9]. In addition, the I–M transition properties in the relaxed films generally depend on the film thickness because the amount of the defects would also depend on it [10]. Since precise control of interface properties plays a decisive role in modern electronics, correlated electron devices also require high quality interfaces where the different types of spin ordering and COO encounter. Therefore, it is desired to establish the technologies to achieve the first-order phase transition in the epitaxial thin films of COO manganites. In this paper, we review the essentials to control the firstorder phase transition arising from the COO in epitaxial thin films, namely, the role of epitaxial strain on (1 1 0) perovskite substrates. The crucial issues in designing the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0) are also discussed, which should be taken into consideration to realize correlated electron devices. In addition, the epitaxial growth of high-Tco materials such as (Bi, Sr)MnO3
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Fig. 1. X-ray diffraction patterns for Nd0.49 Sr0.51 MnO3 : (a) coherent epitaxial and (b) polycrystalline thin films grown on LaAlO3 (0 0 1) substrates. (c) Temperature dependence of resistivity under various magnetic fields of 0, 5, 7, and 9 T for the two types of films.
and A-site ordered SmBaMn2 O6 thin films is presented, aiming at room temperature operation of the devices for a practical use. 2. Experimental Thin films of perovskite manganites, not only narrow and wide one-electron bandwidths COO materials of Pr0.5 Ca0.5 MnO3 [11] and (Nd, Pr, Sr)MnO3 [12,13] but also high-Tco materials of (Bi, Sr)MnO3 [14,15] and SmBaMn2 O6 [16–18], were grown by a pulsedlaser deposition method on various single-crystal substrates of LaAlO3 (LAO) (0 0 1) (a = 3.79 Å), [(LaAlO3 )0.3 –(SrAl0.5 Ta0.5 O3 )0.7 ] (LSAT) (0 0 1), (1 1 0) (a = 3.87 Å), and SrTiO3 (STO) (0 0 1), (1 1 0) (a = 3.905 Å). Growth conditions of substrate temperature and partial oxygen pressure were almost the same as previously reported in the literature [19]. We also tuned these conditions for the growth of (Bi, Sr)MnO3 and SmBaMn2 O6 epitaxial thin films as will be described later. Except for relaxed films, and (Bi, Sr)MnO3 and SmBaMn2 O6 epitaxial thin films, we confirmed layer-by-
layer growth both on (0 0 1) and (1 1 0) substrates monitoring the intensity oscillation of reflection high-energy electron diffraction (RHEED). After the growth, a RHEED pattern consisting of Laue spot and streak was observed, which ensures a smooth surface on a nanometer scale. For these films, atomically flat surfaces with a step and terrace structure were also observed with an atomic force microscopy. The thickness of the thin films is in the range of 60–80 nm except for Nd0.49 Sr0.51 MnO3 and Bi0.5 Sr0.5 MnO3 films (∼300 nm). Four-circle X-ray diffraction measurements were performed to analyze the crystal structure of the films. To confirm the coherent growth of the films, contour mappings in the reciprocal space were also examined around (1 1 4) peak for the films on (0 0 1) substrates, and around (2 2 2) and (3 1 0) peaks for those on (1 1 0) substrates, respectively. A superconducting quantum interference device magnetometer was used to measure the temperature (T) dependence of the magnetization (M) applying a magnetic field of 0.5 T along the in-plane direction during zero-field-cooled fieldwarming (ZFC-FW) and field-cooling (FC) runs. Temperature (T)
Fig. 2. (a) Coordinate axes for (1 1 0) oriented films. Schematics of (b) inclined (∼45◦ ) and (c) perpendicular COO planes to (1 1 0) surfaces. Schematic OO patterns of 3d(x2 − y2 ) anticipated for A-type Pr0.5 Sr0.5 MnO3 thin films on LSAT (1 1 0) substrates (b), and of 3d(3x2 − r2 /3y2 − r2 ) anticipated for CE-type Pr0.5 Ca0.5 MnO3 thin films on LSAT (1 1 0) substrates (c) are also illustrated, respectively. Note that only the OO pattern is correct for the latter, indeed, the COO plane is inclined to the surface of the substrate in this case. Meanwhile, an example of the perpendicular COO plane is reported for LaMnO3 thin films on LSAT (1 1 0) substrates. (d) Shear-mode deformation above (left) and below (right) Too viewed along in-plane [0 0 1] axis.
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dependence of resistivity () under various magnetic fields was evaluated with a standard four-probe method. 3. Results and discussion As described above, a first-order phase transition is hindered in thin films of COO manganites coherently grown on (0 0 1) substrates. For instance, we have already shown that a monotonous temperature dependence of magnetic and transport properties was obtained for Nd0.5 Sr0.5 MnO3 epitaxial thin films on STO (0 0 1) substrates in which tensile strain with a c/a ratio of 0.974 was given [20], while the bulk samples show an FMM phase with Tco of 255 K and an AFMI-COO transition at Tco of 158 K [12]. As another example, a compressive strain effect for Nd0.49 Sr0.51 MnO3 epitaxial films on LAO (0 0 1) substrates is described below. Note that the ground state of Nd0.49 Sr0.51 MnO3 is a CE-type COOI and an A-type two dimensional metallic coexisting state [21]. Fig. 1(a) and (b) displays X-ray diffraction patterns (2–ω scan) for Nd0.49 Sr0.51 MnO3 coherent epitaxial and polycrystalline films comprised of strained and misoriented portions on LAO (0 0 1) substrates, respectively. The former shows only an oriented (0 0 2) peak (2d = 3.93 Å) and the latter also shows other peaks (2d = 3.89, 3.86, and 3.83 Å) corresponding to those in bulk samples due to the lattice relaxation. Incidentally, both films have almost the same thickness of 300 nm, in which the difference between epitaxial and relaxed films was brought about by subtle changes in the deposition process. As for the epitaxial films with a c/a ratio of 1.04, resistivity is highly insulating with a kink at ∼150 K, indicating a possible modification into a C-type OO state. In addition, a negative magneto-resistance under a magnetic field of 9 T was observed as shown in Fig. 1(c). Although the OO planes in compressively strained epitaxial thin films are lying out of the (0 0 1) surface, a first-order phase transition was not observed for the films with the OO planes orthogonal to the (0 0 1) surface. In contrast, partially relaxed films show lower resistivity and an I–M transition recovered by applying a magnetic field of 5 T [Fig. 1(c)], which is due to the disorder effect in relaxed films as pointed out above. These examples clearly prove that both tensile and compressive strain from the (0 0 1) substrates prohibit the firstorder phase transition in epitaxial thin films of COO manganites, and disorder effects are inevitable in relaxed films. To solve the problems above, we proposed the use of (1 1 0) substrates as an essential technology to control the COO in epitaxial thin films [22]. In contrast to the films with a tetragonal symmetry imposed on (0 0 1) substrates, a shear-mode deformation is available for (1 1 0) oriented epitaxial thin films, which allows the anisotropic lattice deformation (J–T q2 mode) necessary for the COO transition without any lattice relaxation. Thus a first-order phase transition can be realized at Too even though the in-plane lattice is fixed to the substrate as shown in Fig. 2(d). The shear-mode deformation for the epitaxial thin films on (1 1 0) substrates was observed together with a discontinuous lattice change at Too and a superlattice reflection corresponding to the OO appeared below Too [23]. Most of the cases, the OO plane is inclined to about 45◦ to the (1 1 0) surface as illustrated in Fig. 2(b), which was also confirmed with the optical anisotropy in transmittance measurements along in-plane [0 0 1] and [1,−1,0] axes [22]. The differences in d(1 1 0) [out-of-plane lattice constant] between (2 2 2) and (3 1 0) reflections in reciprocal space mappings indicate that the films grow inclined to the substrates along [0 0 1] axis with a tilting angle of 0.42◦ and along [1,−1,0] axis with that of 0.31◦ [20], resulting in lower triclinic symmetry [20,24]. The OO plane orthogonal to the (1 1 0) surface as depicted in Fig. 2(c) was observed for LaMnO3 epitaxial thin films grown on LSAT (1 1 0) substrates, because the in-plane d(0 0 1) is the shortest corresponding to the c-axis [the distance between the OO planes] in the pseudo-orthorhombic setting [25].
Fig. 3. Temperature dependence of resistivity for Pr0.5 Ca0.5 MnO3 , Pr0.5 Sr0.5 MnO3 , and Nd0.5 Sr0.5 MnO3 epitaxial thin films coherently grown on LSAT (1 1 0) substrates, together with that for a Pr0.5 Ca0.5 MnO3 /LSAT (0 0 1) sample.
Fig. 3 displays the temperature dependence of resistivity for Pr0.5 Ca0.5 MnO3 , Pr0.5 Sr0.5 MnO3 , and Nd0.5 Sr0.5 MnO3 epitaxial thin films coherently grown on LSAT (1 1 0) substrates. As a control sample, − T for a Pr0.5 Ca0.5 MnO3 thin film on the LSAT (0 0 1) substrate is also shown. As for the Pr0.5 Ca0.5 MnO3 films on the (0 0 1) substrate, the absence of anomaly and the higher values in resistivity are observed. Considering that the favorable lattice distortion for the COO [CE-type, d(3x2 − r2 /3y2 − r2 )] is already given at RT with tetragonal symmetry imposed by the tensile strain on the (0 0 1) substrate, it is natural to anticipate that the OO plane lies in the film plane. Thus we deduce that the (0 0 1) substrate-induced tensile strain will enhance the COO but suppress the first-order phase transition as described above [26]. One can also find that the Pr0.5 Sr0.5 MnO3 film shows a sharp metal–insulator transition while the COO [CE-type, d(3x2 − r2 /3y2 − r2 )] in the Nd0.5 Sr0.5 MnO3 film is modified into a two-phase coexisting state, showing rather metallic behavior on LSAT (1 1 0) substrates [22]. Since the metallic behavior in the Nd0.5 Sr0.5 MnO3 film above is distinguished from the disorder-induced coexisting states in relaxed films, these − T data indicate that the spin and COO can be modulated due to the epitaxial strain. In fact, a one-dimensional elongation of the OO [A-type, d(x2 − y2 )] along [1,−1,0] axis is suggested for Pr0.5 Sr0.5 MnO3 films on LSAT (1 1 0) substrates [23]. As seen from the above, the (1 1 0) epitaxial thin films not always undergo the COO transition similar to the bulk samples. This view is also confirmed by the systematic change in magnetization for the (Nd1−x Prx )0.5 Sr0.5 MnO3 films on LSAT (1 1 0) substrates where x = 0, 1/3, 1/2, and 1. As shown in Fig. 4, a FMM state with thermal hysteresis is gradually turned into the AFM state as a (Pr/Nd) ratio x increases. At x = 1, the Pr0.5 Sr0.5 MnO3 film also shows a sharp FM–AFM transition corresponding to the M–I transition in resistivity measurements. Since such a crossover was not reported in bulk samples, we should also pay attention to the choice of the substrates in order to control the COO. For example, the inverse trend is observed for the (Nd1−x Prx )0.5 Sr0.5 MnO3 films on STO (1 1 0) substrates, hence which should be regarded as a substrate-strain-induced crossover of the COO transition [22]. With the use of (1 1 0) epitaxial technologies, we are now able to go a step further to the atomically sharp interfaces consisting of different types of spin and COO. In fact, recent findings such as superconductivity at the LaAlO3 /SrTiO3 [27] and the quantum Hall effect at the ZnO/MgZnO [28] indeed
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Fig. 5. Cross sectional schematic surface and interface of (1 1 0) oriented films and substrates viewed along [0 0 1] axis. Alternating stacking of (ABO)4+ and O2 4− layers is specific to (1 1 0) surfaces. Conflict of lattice mismatches between in-plane and out-of-pane lattice at step edges, and surface termination with (right) and without O2 layer (left) is also depicted.
Fig. 4. Temperature dependence of magnetization for (Nd1−x Prx )0.5 Sr0.5 MnO3 epitaxial thin films on LSAT (1 1 0) substrates where x = 0, 1/3, 1/2, and 1 (a–d). Thin and bold lines denote the data measured with ZFC-FW and FC runs applying a magnetic field of 0.5 T, respectively.
demonstrated the crucial importance of high quality oxide interfaces. Besides, several intriguing phenomena have hitherto been observed in the superlattices consisting of a Mott insulator and a band insulator such as LaTiO3 /SrTiO3 [29], 3d-transition metal oxides/Nb:SrTiO3 [30,31], and LaMnO3 /SrMnO3 [32,33]. All of the superlattices are fabricated on (0 0 1) substrates, in which the two-dimensional phenomena arising from the charge transfer at the hetero-interfaces are explored. In addition, there are a few studies reporting junction properties in Nd0.5 Sr0.5 MnO3 [34] or La7/8 Sr1/8 MnO3 [35] epitaxial films deposited on the Nb-doped STO (1 1 0) substrates, and an artificially controlled electronic phase separation in Pr0.5 Ca0.5 MnO3 /La0.5 Sr0.5 MnO3 superlattices [36] utilizing the (1 1 0) epitaxial technologies presented in this paper. As a prototypical correlated electron interface, a
(Pr0.5 Sr0.5 MnO3 /Nd0.5 Sr0.5 MnO3 ) interface would be interesting. Prior to fabricating the (Pr0.5 Sr0.5 MnO3 /Nd0.5 Sr0.5 MnO3 ) single interface, we tentatively examined the thickness dependence of the COO to design the number of constituent layers, where the monolayer thickness is defined as d(1 1 0). For the films with a thickness of 23 layers, both Pr0.5 Sr0.5 MnO3 and Nd0.5 Sr0.5 MnO3 ultra thin films hold almost the same properties as described above. Upon reducing the layer thickness down to 11 layers, the FMM phase vanishes due to the two-dimensionality but the COO still survives as evidenced by the thermal hysteresis in a magneto-transport property [37]. In this way, we can choose the number of constituent layers to design the COO interfaces. In addition to the dimensional crossover in ultra thin films, electronic and orbital reconstructions at the surfaces and interfaces should also be crucial as demonstrated by an insulating state due to the oxygen overlayer onto the surface of La5/8 Ca3/8 MnO3 films grown on Nb-doped STO (0 0 1) substrates [38]. However, the problems of surface termination layer on (1 1 0) surfaces and electronic and orbital reconstructions at interfaces are not yet revealed, the situation of which is illustrated in Fig. 5. Focusing on the step edges between films and substrates, a conflict of lattice mismatches between in-plane and out-of-plane lattice constants should exist. In the case of tensile strain, in-plane lattice is elongated and out-of-plane lattice is shrinked; this situation will cause a misfit defect at step edges because the step-height is larger than the out-of-plane lattice constant of the film. Concerning the surface termination layer, the problem of whether the oxygen layer exists or not should be examined with the use of STM in a UHV ambient. Apart from these crucial issues, we can discuss the features inherent to the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0), respectively. At the (0 0 1) interface, the charge transfer is a key issue but a shear-mode deformation cannot be available. In a complementary style, the (1 1 0) interface provides the situation vice versa, if the (1 1 0) interface can be regarded as alternating stackings of nominal charged (ABO)4+ and O2 4− layers. Recently, another view has also been proposed; the amount of the charge transfer is 50% less compared to the case of (0 0 1) interfaces [39]. Thus it is important to determine which view is appropriate to describe the nature of the (1 1 0) interface. A more interesting case is the (2 1 0) interface. Since alternating stackings of AO and BO2 layers similar to those for the (0 0 1) interfaces emerges again, we can expect availability both of charge transfer and a shear-mode deformation at the (2 1 0) interface. We consider that the combinations of (LaMnO3 /SrMnO3 ) and (PrMnO3 /SrMnO3 ) are good candidates to examine the features of the three types of interfaces discussed here. Before closing this section, we briefly describe the epitaxial growth of high-Tco materials towards room temperature operation of correlated electron devices. There are three candidates such as (Bi, Sr)MnO3 [14,15], A-site ordered perovskite manganites [16–18], and bi-layered manganites [Pr(Sr1−y Cay )2 Mn2 O7 (y ≥ 0.8)]
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Fig. 6. Temperature dependence of the phonon mode intensities S (top), frequencies ω (middle), and widths (bottom) in (a) J–T mode and (b) breathing mode, respectively, which were measured in the (x, x) polarization configuration for Bi0.5 Sr0.5 MnO3 epitaxial thin films on STO (0 0 1) substrates.
[40]. As for Bi0.5 Sr0.5 MnO3 compounds, we have already achieved the epitaxial growth of Bi0.5 Sr0.5 MnO3 thin films on STO (0 0 1) substrates by employing the relatively low growth temperature of 480 ◦ C, then conducted Raman spectroscopy measurements to examine the OO [41]. Temperature dependence of the phonon mode intensities S (top), frequencies ω (middle), and widths (bottom) in breathing and J–T modes are shown in Fig. 6(a) and (b), respectively, which were measured in the (x, x) polarization configuration. The result indicates that the OO exists up to 500 K for the films. Meanwhile, we also accomplished the epitaxial growth of SmBaMn2 O6 thin films on STO (0 0 1) substrates. By optimizing the growth conditions with the use of a mixture of Ar and O2 gases, we can obtain high quality epitaxial thin films with an excellent degree of A-site ordering higher than 90% [42]. The next issue we should do is to fabricate these compounds on the (1 1 0) substrates in order to confirm the COO and its phase transition above RT. In fact, Bi0.4 Ca0.6 MnO3 films with a thickness from 300 to 500 nm on STO (1 1 0) and LAO (1 0 0) substrates exhibit a clear COO transition at Tco = 270 K [43], and the thickness dependence was also reported [44]. Note that the films were epitaxial but not coherently grown, namely, the in-plane lattice of the films is not fixed to the substrates. In this case, it is possible even for the films grown on the (0 0 1) substrates to reproduce the COO transition as described in the introduction. However, it is difficult to examine the reason why Tco of the film is lowered. Incidentally, since all of the candidates above are also associated with the multiferroics, we will be able to incorporate the multiferroic and/or ferrotroidic properties as a function of the correlated electron devices. 4. Conclusions We reviewed the essentials to control the first-order phase transition arising from the COO in epitaxial thin films of perovskite manganites, namely, the role of epitaxial strain on (1 1 0) perovskite
substrates. We also pointed out the crucial issues concerning a surface termination layer and a conflict of lattice mismatches at the step edge interface as open questions, and discussed the features of the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0) in terms of charge transfer and a shear-mode deformation as key factors, which should be taken into consideration to design correlated electron devices. The epitaxial growth of high-Tco materials such as (Bi, Sr)MnO3 and A-site ordered SmBaMn2 O6 thin films was briefly described, which is indispensable for aiming at room temperature operation of the devices for a practical use. All of the results and discussions presented here are invaluable in order to establish a basis for realizing correlated electron devices as oxide electronics. Acknowledgements The authors would like to thank Y. Wakabayashi, Y. Uozu, and H. Tamaru for enlightening discussions. References [1] Y. Tomioka, A. Asamitsu, Y. Moritomo, Y. Tokura, J. Phys. Soc. Jpn. 64 (1995) 3626. [2] Y. Tokura, N. Nagaosa, Science 288 (2000) 462. [3] Y. Konishi, Z. Fang, M. Izumi, T. Manako, M. Kasai, H. Kuwahara, M. Kawasaki, K. Terakura, Y. Tokura, J. Phys. Soc. Jpn. 68 (1999) 3790. [4] W. Prellier, A. Biswas, M. Rajeswari, T. Venkatesan, R.L. Green, Appl. Phys. Lett. 75 (1999) 397. [5] H. Oshima, Y. Ishihara, M. Nakamura, K. Miyano, Phys. Rev. B 63 (2000) 094420. [6] S.K. Singh, S.B. Palmer, D.McK. Paul, M.R. Lees, Appl. Phys. Lett. 69 (1996) 263. [7] M. Kasai, H. Kuwahara, Y. Moritomo, Y. Tomioka, Y. Tokura, Jpn. J. Appl. Phys. 35 (1996) L489. [8] M. Kasai, H. Kuwahara, Y. Tomioka, Y. Tokura, J. Appl. Phys. 80 (1996) 6894. [9] Z.Q. Yang, R.W.A. Hendrikx, P.J.M v. Bentum, J. Aarts, Europhys. Lett. 58 (2002) 864. [10] W. Prellier, Ch. Simon, A.M. Haghiri-Gosnet, B. Mercey, B. Raveau, Phys. Rev. B62 (2000) 16337.
56
Y. Ogimoto et al. / Materials Science and Engineering B 173 (2010) 51–56
[11] Y. Tomioka, A. Asasmitsu, H. Kuwahara, Y. Moritomo, Y. Tokura, Phys. Rev. B 53 (1996) 1689. [12] H. Kuwahara, Y. Tomioka, A. Asamitsu, Y. Moritomo, Y. Tokura, Science 270 (1995) 961. [13] H. Kawano, R. Kajimoto, H. Yoshizawa, Y. Tomioka, H. Kuwahara, Y. Tokura, Phys. Rev. Lett. 78 (1997) 4253. ˜ [14] J.L. García-Munoz, C. Frontera, M.A.G. Aranda, A. Llobet, C. Ritter, Phys. Rev. B 63 (2001) 064415. [15] J. Hejtmánek, K. Kníˇzek, Z. Jirák, M. Hervieu, C. Martin, M. Nevˇriva, P. Beran, J. Appl. Phys. 93 (2003) 7370. [16] T. Nakajima, H. Kageyama, Y. Ueda, J. Phys. Chem. Solids 63 (2002) 913. [17] D. Akahoshi, Y. Okimoto, M. Kubota, R. Kumai, T. Arima, Y. Tomioka, Y. Tokura, Phys. Rev. B 70 (2004) 064418. [18] T. Arima, D. Akahoshi, K. Oikawa, T. Kamiyama, M. Uchida, Y. Matsui, Y. Tokura, Phys. Rev. B 66 (2002) 140408. [19] Y. Ogimoto, M. Izumi, T. Manako, T. Kimura, Y. Tomioka, M. Kawasaki, Y. Tokura, Appl. Phys. Lett. 78 (2001) 3505. [20] M. Nakamura, Y. Ogimoto, H. Tamaru, M. Izumi, K. Miyano, Appl. Phys. Lett. 86 (2005) 182504. [21] R. Kajimoto, H. Yoshizawa, H. Kawano, H. Kuwahara, Y. Tokura, K. Ohoyama, M. Ohashi, Phys. Rev. B 60 (1999) 9506. [22] Y. Ogimoto, M. Nakamura, N. Takubo, H. Tamaru, M. Izumi, K. Miyano, Phys. Rev. B 71 (2005) 060403. [23] Y. Wakabayashi, D. Bizen, Y. Kubo, H. Nakao, Y. Murakami, M. Nakamura, Y. Ogimoto, K. Miyano, H. Sawa, J. Phys. Soc. Jpn. 77 (2008) 014712. [24] Y. Ogimoto, N. Takubo, M. Nakamura, H. Tamaru, M. Izumi, K. Miyano, Appl. Phys. Lett. 86 (2005) 112513. [25] H. Tamaru, K. Ishida, N. Ogawa, Y. Kubo, K. Miyano, Phys. Rev. B 78 (2008) 075119. [26] Y. Ogimoto, M. Nakamura, N. Takubo, H. Tamaru, M. Izumi, K. Miyano, Thin Solid Films 486 (2005) 104.
[27] N. Reyren, S. Thiel, A.D. Caviglia, L. Fitting Kourkoutis, G. Hammerl, C. Richter, C.W. Schneider, T. Kopp, A.-S. Rüetschi, D. Jaccard, M. Gabay, D.A. Muller, J.-M. Triscone, J. Mannhart, Science 317 (2007) 1196. [28] A. Tsukazaki, A. Ohtomo, T. Kita, Y. Ohno, H. Ohno, M. Kawasaki, Science 315 (2007) 1388. [29] A. Ohtomo, D.A. Muller, J.L. Grazul, H.Y. Hwang, Nature 419 (2002) 378. [30] A. Sawa, A. Yamamoto, H. Yamada, T. Fujii, M. Kawasaki, J. Matsuno, Y. Tokura, Appl. Phys. Lett. 90 (2007) 252102. [31] M. Nakamura, A. Sawa, H. Sato, H. Akoh, M. Kawasaki, Y. Tokura, Phys. Rev. B 75 (2007) 155103. [32] H. Yamada, M. Kawasaki, T. Lottermoser, T. Arima, Y. Tokura, Appl. Phys. Lett. 89 (2006) 052506. [33] N. Ogawa, T. Satoh, Y. Ogimoto, K. Miyano, Phys. Rev. B 78 (2008) 212409. [34] J. Matsuno, A. Sawa, M. Kawasaki, Y. Tokura, Appl. Phys. Lett. 92 (2008) 122104. [35] Y.Z. Chen, J.R. Sun, A.D. Wei, W.M. Lu, S. Liang, B.G. Shen, Appl. Phys. Lett. 93 (2008) 152515. [36] M. Nakamura, D. Okuyama, J.S. Lee, T. Arima, Y. Wakabayashi, R. Kumai, M. Kawasaki, Y. Tokura, Adv. Mater. 21 (2009) 1. [37] N. Harada, Y. Ogimoto, N. Ogawa, K. Miyano, unpublished. [38] K. Fuchigami, Z. Gai, T.Z. Ward, L.F. Yin, P.C. Snijders, E.W. Plummer, J. Shen, Phys. Rev. Lett. 102 (2009) 066104. [39] J.X. Ma, X.F. Liu, T. Lin, G.Y. Gao, J.P. Zhang, W.B. Wu, X.G. Li, J. Shi, Phys. Rev. B 79 (2009) 174424. [40] Y. Tokunaga, T.J. Sato, M. Uchida, R. Kumai, Y. Matsui, T. Arima, Y. Tokura, Phys. Rev. B 77 (2008) 064428. [41] T. Kawasaki, Y. Ogimoto, N. Ogawa, K. Miyano, H. Tamaru, M. Izumi, J. Appl. Phys. 101 (2007) 123714. [42] Y. Kubo, Y. Ogimoto, K. Miyano, in preparation. [43] D.H. Kim, H.M. Christen, M. Varela, H.N. Lee, D.H. Lowndes, Appl. Phys. Lett. 88 (2008) 202503. [44] Y.Z. Chen, J.R. Sun, S. Liang, W.M. Lv, B.G. Shen, W.B. Wu, J. Appl. Phys. 103 (2008) 096105.
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Materials Science and Engineering B journal homepage: www.elsevier.com/locate/mseb
The essentials for the control of charge-orbital ordering in thin films of perovskite manganites Yasushi Ogimoto a,b,c,∗ , Masao Nakamura d,1 , Nobutaka Harada d , Naoki Ogawa a,b , Kenjiro Miyano a,b a
Research Center for Advanced Science and Technology (RCAST), The University of Tokyo, Tokyo 153-8904, Japan CREST, Japan Science and Technology Agency, Saitama 332-0012, Japan Fuji Electric Advanced Technology Co., Ltd., Tokyo 191-8502, Japan d Department of Applied Physics, The University of Tokyo, Tokyo 113-8656, Japan b c
a r t i c l e
i n f o
Article history: Received 1 July 2009 Received in revised form 1 February 2010 Accepted 3 February 2010 Keywords: Charge-orbital ordering Phase transition Epitaxial strain Correlated electron interface
a b s t r a c t Gigantic responses arising from the melting of charge-orbital ordering (COO) in perovskite manganites have been attractive both in material science and technologies. We review the essentials to realize such phenomena in a thin film form aiming at a possible application to correlated electron devices, namely, the role of epitaxial strain on (1 1 0) substrates. The crucial issues in designing the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0) are also discussed. In addition, the epitaxial growth of high-Tco materials such as (Bi, Sr)MnO3 and A-site ordered SmBaMn2 O6 thin films is described, which is indispensable towards room temperature operation of correlated electron devices. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Gigantic and ultrafast responses arising from the melting of charge-orbital ordering (COO) in perovskite manganites have been attractive both in material science and technologies, because of the rich physics underlying an antiferromagnetic insulator (AFMI) to a ferromagnetic metal (FMM) transition and its possible applicability to correlated electron devices [1,2]. The transition between the FMM and the COO states is first-order, and is coupled to the lattice degree of freedom through the Jahn–Teller (J–T) interaction. Therefore, it is of crucial importance to have a deep understanding of an epitaxial strain from the substrates to control the COO and its AFMI–FMM transition in a thin film form. In the case of the films coherently grown on (0 0 1) perovskite substrates, Konishi et al. demonstrated that the ground states of La0.5 Sr0.5 MnO3 epitaxial thin films are drastically modulated by the substrate-induced strain through the OO close to the phase boundaries adjoining Atype AFMI, FMM, and C-type AFMI states [3]. This striking example also means that the first-order phase transition associated with the COO is prohibited for the epitaxial films coherently grown on (0 0 1)
∗ Corresponding author at: Research Center for Advanced Science and Technology (RCAST), The University of Tokyo, Photonics Materials, 4-6-1, Komaba, Meguro-ku, Tokyo 153-8904, Japan. Tel.: +81 3 5452 5076; fax: +81 3 5452 5076. E-mail addresses: y [email protected], y [email protected] (Y. Ogimoto). 1 Present address: Cross-correlated Materials Research Group (CMRG), RIKEN, Saitama 351-0198, Japan. 0921-5107/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2010.02.009
substrates. The difficulty arises from severe restriction on the lattice deformation in which the in-plane lattice is clamped to the substrate, and the type of which is limited to an orthogonal-mode due to the tetragonal symmetry on (0 0 1) substrates [4]. To escape from the strain effect above, it is easy to make a relaxed film utilizing a large lattice mismatch between the films and the substrates [5]. Since the in-plane lattice is not restricted to the substrates in the relaxed films, the lattice deformation associated with the I–M transition is possible similar to that of the bulk sample, which could lead to the reproduction of the bulk properties [5,6]. However, the films are polycrystalline with a large amount of defects (e.g., misfit dislocations and grain boundaries) which will make the I–M transition diffusive due to the disorder effect [7–9]. In addition, the I–M transition properties in the relaxed films generally depend on the film thickness because the amount of the defects would also depend on it [10]. Since precise control of interface properties plays a decisive role in modern electronics, correlated electron devices also require high quality interfaces where the different types of spin ordering and COO encounter. Therefore, it is desired to establish the technologies to achieve the first-order phase transition in the epitaxial thin films of COO manganites. In this paper, we review the essentials to control the firstorder phase transition arising from the COO in epitaxial thin films, namely, the role of epitaxial strain on (1 1 0) perovskite substrates. The crucial issues in designing the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0) are also discussed, which should be taken into consideration to realize correlated electron devices. In addition, the epitaxial growth of high-Tco materials such as (Bi, Sr)MnO3
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Fig. 1. X-ray diffraction patterns for Nd0.49 Sr0.51 MnO3 : (a) coherent epitaxial and (b) polycrystalline thin films grown on LaAlO3 (0 0 1) substrates. (c) Temperature dependence of resistivity under various magnetic fields of 0, 5, 7, and 9 T for the two types of films.
and A-site ordered SmBaMn2 O6 thin films is presented, aiming at room temperature operation of the devices for a practical use. 2. Experimental Thin films of perovskite manganites, not only narrow and wide one-electron bandwidths COO materials of Pr0.5 Ca0.5 MnO3 [11] and (Nd, Pr, Sr)MnO3 [12,13] but also high-Tco materials of (Bi, Sr)MnO3 [14,15] and SmBaMn2 O6 [16–18], were grown by a pulsedlaser deposition method on various single-crystal substrates of LaAlO3 (LAO) (0 0 1) (a = 3.79 Å), [(LaAlO3 )0.3 –(SrAl0.5 Ta0.5 O3 )0.7 ] (LSAT) (0 0 1), (1 1 0) (a = 3.87 Å), and SrTiO3 (STO) (0 0 1), (1 1 0) (a = 3.905 Å). Growth conditions of substrate temperature and partial oxygen pressure were almost the same as previously reported in the literature [19]. We also tuned these conditions for the growth of (Bi, Sr)MnO3 and SmBaMn2 O6 epitaxial thin films as will be described later. Except for relaxed films, and (Bi, Sr)MnO3 and SmBaMn2 O6 epitaxial thin films, we confirmed layer-by-
layer growth both on (0 0 1) and (1 1 0) substrates monitoring the intensity oscillation of reflection high-energy electron diffraction (RHEED). After the growth, a RHEED pattern consisting of Laue spot and streak was observed, which ensures a smooth surface on a nanometer scale. For these films, atomically flat surfaces with a step and terrace structure were also observed with an atomic force microscopy. The thickness of the thin films is in the range of 60–80 nm except for Nd0.49 Sr0.51 MnO3 and Bi0.5 Sr0.5 MnO3 films (∼300 nm). Four-circle X-ray diffraction measurements were performed to analyze the crystal structure of the films. To confirm the coherent growth of the films, contour mappings in the reciprocal space were also examined around (1 1 4) peak for the films on (0 0 1) substrates, and around (2 2 2) and (3 1 0) peaks for those on (1 1 0) substrates, respectively. A superconducting quantum interference device magnetometer was used to measure the temperature (T) dependence of the magnetization (M) applying a magnetic field of 0.5 T along the in-plane direction during zero-field-cooled fieldwarming (ZFC-FW) and field-cooling (FC) runs. Temperature (T)
Fig. 2. (a) Coordinate axes for (1 1 0) oriented films. Schematics of (b) inclined (∼45◦ ) and (c) perpendicular COO planes to (1 1 0) surfaces. Schematic OO patterns of 3d(x2 − y2 ) anticipated for A-type Pr0.5 Sr0.5 MnO3 thin films on LSAT (1 1 0) substrates (b), and of 3d(3x2 − r2 /3y2 − r2 ) anticipated for CE-type Pr0.5 Ca0.5 MnO3 thin films on LSAT (1 1 0) substrates (c) are also illustrated, respectively. Note that only the OO pattern is correct for the latter, indeed, the COO plane is inclined to the surface of the substrate in this case. Meanwhile, an example of the perpendicular COO plane is reported for LaMnO3 thin films on LSAT (1 1 0) substrates. (d) Shear-mode deformation above (left) and below (right) Too viewed along in-plane [0 0 1] axis.
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dependence of resistivity () under various magnetic fields was evaluated with a standard four-probe method. 3. Results and discussion As described above, a first-order phase transition is hindered in thin films of COO manganites coherently grown on (0 0 1) substrates. For instance, we have already shown that a monotonous temperature dependence of magnetic and transport properties was obtained for Nd0.5 Sr0.5 MnO3 epitaxial thin films on STO (0 0 1) substrates in which tensile strain with a c/a ratio of 0.974 was given [20], while the bulk samples show an FMM phase with Tco of 255 K and an AFMI-COO transition at Tco of 158 K [12]. As another example, a compressive strain effect for Nd0.49 Sr0.51 MnO3 epitaxial films on LAO (0 0 1) substrates is described below. Note that the ground state of Nd0.49 Sr0.51 MnO3 is a CE-type COOI and an A-type two dimensional metallic coexisting state [21]. Fig. 1(a) and (b) displays X-ray diffraction patterns (2–ω scan) for Nd0.49 Sr0.51 MnO3 coherent epitaxial and polycrystalline films comprised of strained and misoriented portions on LAO (0 0 1) substrates, respectively. The former shows only an oriented (0 0 2) peak (2d = 3.93 Å) and the latter also shows other peaks (2d = 3.89, 3.86, and 3.83 Å) corresponding to those in bulk samples due to the lattice relaxation. Incidentally, both films have almost the same thickness of 300 nm, in which the difference between epitaxial and relaxed films was brought about by subtle changes in the deposition process. As for the epitaxial films with a c/a ratio of 1.04, resistivity is highly insulating with a kink at ∼150 K, indicating a possible modification into a C-type OO state. In addition, a negative magneto-resistance under a magnetic field of 9 T was observed as shown in Fig. 1(c). Although the OO planes in compressively strained epitaxial thin films are lying out of the (0 0 1) surface, a first-order phase transition was not observed for the films with the OO planes orthogonal to the (0 0 1) surface. In contrast, partially relaxed films show lower resistivity and an I–M transition recovered by applying a magnetic field of 5 T [Fig. 1(c)], which is due to the disorder effect in relaxed films as pointed out above. These examples clearly prove that both tensile and compressive strain from the (0 0 1) substrates prohibit the firstorder phase transition in epitaxial thin films of COO manganites, and disorder effects are inevitable in relaxed films. To solve the problems above, we proposed the use of (1 1 0) substrates as an essential technology to control the COO in epitaxial thin films [22]. In contrast to the films with a tetragonal symmetry imposed on (0 0 1) substrates, a shear-mode deformation is available for (1 1 0) oriented epitaxial thin films, which allows the anisotropic lattice deformation (J–T q2 mode) necessary for the COO transition without any lattice relaxation. Thus a first-order phase transition can be realized at Too even though the in-plane lattice is fixed to the substrate as shown in Fig. 2(d). The shear-mode deformation for the epitaxial thin films on (1 1 0) substrates was observed together with a discontinuous lattice change at Too and a superlattice reflection corresponding to the OO appeared below Too [23]. Most of the cases, the OO plane is inclined to about 45◦ to the (1 1 0) surface as illustrated in Fig. 2(b), which was also confirmed with the optical anisotropy in transmittance measurements along in-plane [0 0 1] and [1,−1,0] axes [22]. The differences in d(1 1 0) [out-of-plane lattice constant] between (2 2 2) and (3 1 0) reflections in reciprocal space mappings indicate that the films grow inclined to the substrates along [0 0 1] axis with a tilting angle of 0.42◦ and along [1,−1,0] axis with that of 0.31◦ [20], resulting in lower triclinic symmetry [20,24]. The OO plane orthogonal to the (1 1 0) surface as depicted in Fig. 2(c) was observed for LaMnO3 epitaxial thin films grown on LSAT (1 1 0) substrates, because the in-plane d(0 0 1) is the shortest corresponding to the c-axis [the distance between the OO planes] in the pseudo-orthorhombic setting [25].
Fig. 3. Temperature dependence of resistivity for Pr0.5 Ca0.5 MnO3 , Pr0.5 Sr0.5 MnO3 , and Nd0.5 Sr0.5 MnO3 epitaxial thin films coherently grown on LSAT (1 1 0) substrates, together with that for a Pr0.5 Ca0.5 MnO3 /LSAT (0 0 1) sample.
Fig. 3 displays the temperature dependence of resistivity for Pr0.5 Ca0.5 MnO3 , Pr0.5 Sr0.5 MnO3 , and Nd0.5 Sr0.5 MnO3 epitaxial thin films coherently grown on LSAT (1 1 0) substrates. As a control sample, − T for a Pr0.5 Ca0.5 MnO3 thin film on the LSAT (0 0 1) substrate is also shown. As for the Pr0.5 Ca0.5 MnO3 films on the (0 0 1) substrate, the absence of anomaly and the higher values in resistivity are observed. Considering that the favorable lattice distortion for the COO [CE-type, d(3x2 − r2 /3y2 − r2 )] is already given at RT with tetragonal symmetry imposed by the tensile strain on the (0 0 1) substrate, it is natural to anticipate that the OO plane lies in the film plane. Thus we deduce that the (0 0 1) substrate-induced tensile strain will enhance the COO but suppress the first-order phase transition as described above [26]. One can also find that the Pr0.5 Sr0.5 MnO3 film shows a sharp metal–insulator transition while the COO [CE-type, d(3x2 − r2 /3y2 − r2 )] in the Nd0.5 Sr0.5 MnO3 film is modified into a two-phase coexisting state, showing rather metallic behavior on LSAT (1 1 0) substrates [22]. Since the metallic behavior in the Nd0.5 Sr0.5 MnO3 film above is distinguished from the disorder-induced coexisting states in relaxed films, these − T data indicate that the spin and COO can be modulated due to the epitaxial strain. In fact, a one-dimensional elongation of the OO [A-type, d(x2 − y2 )] along [1,−1,0] axis is suggested for Pr0.5 Sr0.5 MnO3 films on LSAT (1 1 0) substrates [23]. As seen from the above, the (1 1 0) epitaxial thin films not always undergo the COO transition similar to the bulk samples. This view is also confirmed by the systematic change in magnetization for the (Nd1−x Prx )0.5 Sr0.5 MnO3 films on LSAT (1 1 0) substrates where x = 0, 1/3, 1/2, and 1. As shown in Fig. 4, a FMM state with thermal hysteresis is gradually turned into the AFM state as a (Pr/Nd) ratio x increases. At x = 1, the Pr0.5 Sr0.5 MnO3 film also shows a sharp FM–AFM transition corresponding to the M–I transition in resistivity measurements. Since such a crossover was not reported in bulk samples, we should also pay attention to the choice of the substrates in order to control the COO. For example, the inverse trend is observed for the (Nd1−x Prx )0.5 Sr0.5 MnO3 films on STO (1 1 0) substrates, hence which should be regarded as a substrate-strain-induced crossover of the COO transition [22]. With the use of (1 1 0) epitaxial technologies, we are now able to go a step further to the atomically sharp interfaces consisting of different types of spin and COO. In fact, recent findings such as superconductivity at the LaAlO3 /SrTiO3 [27] and the quantum Hall effect at the ZnO/MgZnO [28] indeed
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Fig. 5. Cross sectional schematic surface and interface of (1 1 0) oriented films and substrates viewed along [0 0 1] axis. Alternating stacking of (ABO)4+ and O2 4− layers is specific to (1 1 0) surfaces. Conflict of lattice mismatches between in-plane and out-of-pane lattice at step edges, and surface termination with (right) and without O2 layer (left) is also depicted.
Fig. 4. Temperature dependence of magnetization for (Nd1−x Prx )0.5 Sr0.5 MnO3 epitaxial thin films on LSAT (1 1 0) substrates where x = 0, 1/3, 1/2, and 1 (a–d). Thin and bold lines denote the data measured with ZFC-FW and FC runs applying a magnetic field of 0.5 T, respectively.
demonstrated the crucial importance of high quality oxide interfaces. Besides, several intriguing phenomena have hitherto been observed in the superlattices consisting of a Mott insulator and a band insulator such as LaTiO3 /SrTiO3 [29], 3d-transition metal oxides/Nb:SrTiO3 [30,31], and LaMnO3 /SrMnO3 [32,33]. All of the superlattices are fabricated on (0 0 1) substrates, in which the two-dimensional phenomena arising from the charge transfer at the hetero-interfaces are explored. In addition, there are a few studies reporting junction properties in Nd0.5 Sr0.5 MnO3 [34] or La7/8 Sr1/8 MnO3 [35] epitaxial films deposited on the Nb-doped STO (1 1 0) substrates, and an artificially controlled electronic phase separation in Pr0.5 Ca0.5 MnO3 /La0.5 Sr0.5 MnO3 superlattices [36] utilizing the (1 1 0) epitaxial technologies presented in this paper. As a prototypical correlated electron interface, a
(Pr0.5 Sr0.5 MnO3 /Nd0.5 Sr0.5 MnO3 ) interface would be interesting. Prior to fabricating the (Pr0.5 Sr0.5 MnO3 /Nd0.5 Sr0.5 MnO3 ) single interface, we tentatively examined the thickness dependence of the COO to design the number of constituent layers, where the monolayer thickness is defined as d(1 1 0). For the films with a thickness of 23 layers, both Pr0.5 Sr0.5 MnO3 and Nd0.5 Sr0.5 MnO3 ultra thin films hold almost the same properties as described above. Upon reducing the layer thickness down to 11 layers, the FMM phase vanishes due to the two-dimensionality but the COO still survives as evidenced by the thermal hysteresis in a magneto-transport property [37]. In this way, we can choose the number of constituent layers to design the COO interfaces. In addition to the dimensional crossover in ultra thin films, electronic and orbital reconstructions at the surfaces and interfaces should also be crucial as demonstrated by an insulating state due to the oxygen overlayer onto the surface of La5/8 Ca3/8 MnO3 films grown on Nb-doped STO (0 0 1) substrates [38]. However, the problems of surface termination layer on (1 1 0) surfaces and electronic and orbital reconstructions at interfaces are not yet revealed, the situation of which is illustrated in Fig. 5. Focusing on the step edges between films and substrates, a conflict of lattice mismatches between in-plane and out-of-plane lattice constants should exist. In the case of tensile strain, in-plane lattice is elongated and out-of-plane lattice is shrinked; this situation will cause a misfit defect at step edges because the step-height is larger than the out-of-plane lattice constant of the film. Concerning the surface termination layer, the problem of whether the oxygen layer exists or not should be examined with the use of STM in a UHV ambient. Apart from these crucial issues, we can discuss the features inherent to the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0), respectively. At the (0 0 1) interface, the charge transfer is a key issue but a shear-mode deformation cannot be available. In a complementary style, the (1 1 0) interface provides the situation vice versa, if the (1 1 0) interface can be regarded as alternating stackings of nominal charged (ABO)4+ and O2 4− layers. Recently, another view has also been proposed; the amount of the charge transfer is 50% less compared to the case of (0 0 1) interfaces [39]. Thus it is important to determine which view is appropriate to describe the nature of the (1 1 0) interface. A more interesting case is the (2 1 0) interface. Since alternating stackings of AO and BO2 layers similar to those for the (0 0 1) interfaces emerges again, we can expect availability both of charge transfer and a shear-mode deformation at the (2 1 0) interface. We consider that the combinations of (LaMnO3 /SrMnO3 ) and (PrMnO3 /SrMnO3 ) are good candidates to examine the features of the three types of interfaces discussed here. Before closing this section, we briefly describe the epitaxial growth of high-Tco materials towards room temperature operation of correlated electron devices. There are three candidates such as (Bi, Sr)MnO3 [14,15], A-site ordered perovskite manganites [16–18], and bi-layered manganites [Pr(Sr1−y Cay )2 Mn2 O7 (y ≥ 0.8)]
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Fig. 6. Temperature dependence of the phonon mode intensities S (top), frequencies ω (middle), and widths (bottom) in (a) J–T mode and (b) breathing mode, respectively, which were measured in the (x, x) polarization configuration for Bi0.5 Sr0.5 MnO3 epitaxial thin films on STO (0 0 1) substrates.
[40]. As for Bi0.5 Sr0.5 MnO3 compounds, we have already achieved the epitaxial growth of Bi0.5 Sr0.5 MnO3 thin films on STO (0 0 1) substrates by employing the relatively low growth temperature of 480 ◦ C, then conducted Raman spectroscopy measurements to examine the OO [41]. Temperature dependence of the phonon mode intensities S (top), frequencies ω (middle), and widths (bottom) in breathing and J–T modes are shown in Fig. 6(a) and (b), respectively, which were measured in the (x, x) polarization configuration. The result indicates that the OO exists up to 500 K for the films. Meanwhile, we also accomplished the epitaxial growth of SmBaMn2 O6 thin films on STO (0 0 1) substrates. By optimizing the growth conditions with the use of a mixture of Ar and O2 gases, we can obtain high quality epitaxial thin films with an excellent degree of A-site ordering higher than 90% [42]. The next issue we should do is to fabricate these compounds on the (1 1 0) substrates in order to confirm the COO and its phase transition above RT. In fact, Bi0.4 Ca0.6 MnO3 films with a thickness from 300 to 500 nm on STO (1 1 0) and LAO (1 0 0) substrates exhibit a clear COO transition at Tco = 270 K [43], and the thickness dependence was also reported [44]. Note that the films were epitaxial but not coherently grown, namely, the in-plane lattice of the films is not fixed to the substrates. In this case, it is possible even for the films grown on the (0 0 1) substrates to reproduce the COO transition as described in the introduction. However, it is difficult to examine the reason why Tco of the film is lowered. Incidentally, since all of the candidates above are also associated with the multiferroics, we will be able to incorporate the multiferroic and/or ferrotroidic properties as a function of the correlated electron devices. 4. Conclusions We reviewed the essentials to control the first-order phase transition arising from the COO in epitaxial thin films of perovskite manganites, namely, the role of epitaxial strain on (1 1 0) perovskite
substrates. We also pointed out the crucial issues concerning a surface termination layer and a conflict of lattice mismatches at the step edge interface as open questions, and discussed the features of the correlated electron interfaces of (0 0 1), (1 1 0), and (2 1 0) in terms of charge transfer and a shear-mode deformation as key factors, which should be taken into consideration to design correlated electron devices. The epitaxial growth of high-Tco materials such as (Bi, Sr)MnO3 and A-site ordered SmBaMn2 O6 thin films was briefly described, which is indispensable for aiming at room temperature operation of the devices for a practical use. All of the results and discussions presented here are invaluable in order to establish a basis for realizing correlated electron devices as oxide electronics. Acknowledgements The authors would like to thank Y. Wakabayashi, Y. Uozu, and H. Tamaru for enlightening discussions. References [1] Y. Tomioka, A. Asamitsu, Y. Moritomo, Y. Tokura, J. Phys. Soc. Jpn. 64 (1995) 3626. [2] Y. Tokura, N. Nagaosa, Science 288 (2000) 462. [3] Y. Konishi, Z. Fang, M. Izumi, T. Manako, M. Kasai, H. Kuwahara, M. Kawasaki, K. Terakura, Y. Tokura, J. Phys. Soc. Jpn. 68 (1999) 3790. [4] W. Prellier, A. Biswas, M. Rajeswari, T. Venkatesan, R.L. Green, Appl. Phys. Lett. 75 (1999) 397. [5] H. Oshima, Y. Ishihara, M. Nakamura, K. Miyano, Phys. Rev. B 63 (2000) 094420. [6] S.K. Singh, S.B. Palmer, D.McK. Paul, M.R. Lees, Appl. Phys. Lett. 69 (1996) 263. [7] M. Kasai, H. Kuwahara, Y. Moritomo, Y. Tomioka, Y. Tokura, Jpn. J. Appl. Phys. 35 (1996) L489. [8] M. Kasai, H. Kuwahara, Y. Tomioka, Y. Tokura, J. Appl. Phys. 80 (1996) 6894. [9] Z.Q. Yang, R.W.A. Hendrikx, P.J.M v. Bentum, J. Aarts, Europhys. Lett. 58 (2002) 864. [10] W. Prellier, Ch. Simon, A.M. Haghiri-Gosnet, B. Mercey, B. Raveau, Phys. Rev. B62 (2000) 16337.
56
Y. Ogimoto et al. / Materials Science and Engineering B 173 (2010) 51–56
[11] Y. Tomioka, A. Asasmitsu, H. Kuwahara, Y. Moritomo, Y. Tokura, Phys. Rev. B 53 (1996) 1689. [12] H. Kuwahara, Y. Tomioka, A. Asamitsu, Y. Moritomo, Y. Tokura, Science 270 (1995) 961. [13] H. Kawano, R. Kajimoto, H. Yoshizawa, Y. Tomioka, H. Kuwahara, Y. Tokura, Phys. Rev. Lett. 78 (1997) 4253. ˜ [14] J.L. García-Munoz, C. Frontera, M.A.G. Aranda, A. Llobet, C. Ritter, Phys. Rev. B 63 (2001) 064415. [15] J. Hejtmánek, K. Kníˇzek, Z. Jirák, M. Hervieu, C. Martin, M. Nevˇriva, P. Beran, J. Appl. Phys. 93 (2003) 7370. [16] T. Nakajima, H. Kageyama, Y. Ueda, J. Phys. Chem. Solids 63 (2002) 913. [17] D. Akahoshi, Y. Okimoto, M. Kubota, R. Kumai, T. Arima, Y. Tomioka, Y. Tokura, Phys. Rev. B 70 (2004) 064418. [18] T. Arima, D. Akahoshi, K. Oikawa, T. Kamiyama, M. Uchida, Y. Matsui, Y. Tokura, Phys. Rev. B 66 (2002) 140408. [19] Y. Ogimoto, M. Izumi, T. Manako, T. Kimura, Y. Tomioka, M. Kawasaki, Y. Tokura, Appl. Phys. Lett. 78 (2001) 3505. [20] M. Nakamura, Y. Ogimoto, H. Tamaru, M. Izumi, K. Miyano, Appl. Phys. Lett. 86 (2005) 182504. [21] R. Kajimoto, H. Yoshizawa, H. Kawano, H. Kuwahara, Y. Tokura, K. Ohoyama, M. Ohashi, Phys. Rev. B 60 (1999) 9506. [22] Y. Ogimoto, M. Nakamura, N. Takubo, H. Tamaru, M. Izumi, K. Miyano, Phys. Rev. B 71 (2005) 060403. [23] Y. Wakabayashi, D. Bizen, Y. Kubo, H. Nakao, Y. Murakami, M. Nakamura, Y. Ogimoto, K. Miyano, H. Sawa, J. Phys. Soc. Jpn. 77 (2008) 014712. [24] Y. Ogimoto, N. Takubo, M. Nakamura, H. Tamaru, M. Izumi, K. Miyano, Appl. Phys. Lett. 86 (2005) 112513. [25] H. Tamaru, K. Ishida, N. Ogawa, Y. Kubo, K. Miyano, Phys. Rev. B 78 (2008) 075119. [26] Y. Ogimoto, M. Nakamura, N. Takubo, H. Tamaru, M. Izumi, K. Miyano, Thin Solid Films 486 (2005) 104.
[27] N. Reyren, S. Thiel, A.D. Caviglia, L. Fitting Kourkoutis, G. Hammerl, C. Richter, C.W. Schneider, T. Kopp, A.-S. Rüetschi, D. Jaccard, M. Gabay, D.A. Muller, J.-M. Triscone, J. Mannhart, Science 317 (2007) 1196. [28] A. Tsukazaki, A. Ohtomo, T. Kita, Y. Ohno, H. Ohno, M. Kawasaki, Science 315 (2007) 1388. [29] A. Ohtomo, D.A. Muller, J.L. Grazul, H.Y. Hwang, Nature 419 (2002) 378. [30] A. Sawa, A. Yamamoto, H. Yamada, T. Fujii, M. Kawasaki, J. Matsuno, Y. Tokura, Appl. Phys. Lett. 90 (2007) 252102. [31] M. Nakamura, A. Sawa, H. Sato, H. Akoh, M. Kawasaki, Y. Tokura, Phys. Rev. B 75 (2007) 155103. [32] H. Yamada, M. Kawasaki, T. Lottermoser, T. Arima, Y. Tokura, Appl. Phys. Lett. 89 (2006) 052506. [33] N. Ogawa, T. Satoh, Y. Ogimoto, K. Miyano, Phys. Rev. B 78 (2008) 212409. [34] J. Matsuno, A. Sawa, M. Kawasaki, Y. Tokura, Appl. Phys. Lett. 92 (2008) 122104. [35] Y.Z. Chen, J.R. Sun, A.D. Wei, W.M. Lu, S. Liang, B.G. Shen, Appl. Phys. Lett. 93 (2008) 152515. [36] M. Nakamura, D. Okuyama, J.S. Lee, T. Arima, Y. Wakabayashi, R. Kumai, M. Kawasaki, Y. Tokura, Adv. Mater. 21 (2009) 1. [37] N. Harada, Y. Ogimoto, N. Ogawa, K. Miyano, unpublished. [38] K. Fuchigami, Z. Gai, T.Z. Ward, L.F. Yin, P.C. Snijders, E.W. Plummer, J. Shen, Phys. Rev. Lett. 102 (2009) 066104. [39] J.X. Ma, X.F. Liu, T. Lin, G.Y. Gao, J.P. Zhang, W.B. Wu, X.G. Li, J. Shi, Phys. Rev. B 79 (2009) 174424. [40] Y. Tokunaga, T.J. Sato, M. Uchida, R. Kumai, Y. Matsui, T. Arima, Y. Tokura, Phys. Rev. B 77 (2008) 064428. [41] T. Kawasaki, Y. Ogimoto, N. Ogawa, K. Miyano, H. Tamaru, M. Izumi, J. Appl. Phys. 101 (2007) 123714. [42] Y. Kubo, Y. Ogimoto, K. Miyano, in preparation. [43] D.H. Kim, H.M. Christen, M. Varela, H.N. Lee, D.H. Lowndes, Appl. Phys. Lett. 88 (2008) 202503. [44] Y.Z. Chen, J.R. Sun, S. Liang, W.M. Lv, B.G. Shen, W.B. Wu, J. Appl. Phys. 103 (2008) 096105.