3)O3–PbTiO3 thin films by using miscut substrates

Thin Solid Films 589 (2015) 792–797

Contents lists available at ScienceDirect

Thin Solid Films journal homepage: www.elsevier.com/locate/tsf

Tuning structure in epitaxial Pb(Mg1/3Nb2/3)O3–PbTiO3 thin films by using miscut substrates M. Mietschke a,b,⁎, S. Oswald a, S. Fähler a, L. Schultz a,b, R. Hühne a a b

IFW Dresden, P.O. Box 270116, D-01171 Dresden, Germany Dresden University of Technology, Faculty of Mechanical Science and Engineering, D-01062 Dresden, Germany

a r t i c l e

i n f o

Article history: Received 17 April 2015 Received in revised form 4 June 2015 Accepted 19 June 2015 Available online 26 June 2015 Keywords: PMN–PT Epitaxial growth Miscut films Ferroelectric films

a b s t r a c t Pb(Mg1/3Nb2/3)O3–PbTiO3 (PMN–PT) is one of the most promising ferroelectric material for actuator, dielectric and electrocaloric applications. However, oriented and phase pure thin films are essential to use the outstanding properties of these compounds. In this work it is demonstrated that the use of miscut substrates influences the growth mechanism leading to a significantly broader deposition window to achieve the required film quality. Therefore, epitaxial 0.68Pb(Mg1/3Nb2/3)O3–0.32PbTiO3 films were grown by pulsed laser deposition on (001)oriented single crystalline SrTiO3 (STO) substrates with a miscut angle between 0 and 15° towards the [100] direction using a conducting La0.7Sr0.3CoO3 buffer layer. The influence of the vicinal angle on the PMN–PT structure was studied by high resolution X-ray diffraction, X-ray photoelectron spectroscopy and transmission electron microscopy. A nearly pure perovskite phase growth with a cube-on-cube epitaxial relationship was obtained on all miscut STO substrates, whereas a significant volume fraction of the pyrochlore phase was present on the standard substrate. Reciprocal space measurements revealed a peak split of the perovskite reflections indicating structural variants of PMN–PT with different c/a ratios. An additional tilting of the PMN–PT planes with respect to the buffer layer was observed on some samples, which might be explained with the incorporation of dislocations according to the Nagai model. Polarization loops were measured in a temperature range between room temperature and 150 °C showing a sharp drop of the remanent polarization above 65 °C on vicinal substrates. © 2015 Elsevier B.V. All rights reserved.

1. Introduction The relaxor based ferroelectric oxide (1 − x)Pb(Mg1/3Nb2/3)O3– xPbTiO3 (PMN–PT) shows a giant electrocaloric effect (ECE) [1] in addition to other functional properties as large dielectric constants and ultrahigh piezoelectric coefficients [2–5]. Such an ECE, which is induced by the application of an electrical field at a temperature close to the diffusionless phase transition, might be used for novel solid state cooling devices [6]. Whereas the typical temperature change is relatively low in bulk materials mainly due to the low applicable electric fields [7], it is much easier to apply higher fields in thin films resulting in a larger ECE for this configuration [8]. Our general interest is to study the correlation between the local microstructure, the ferroelectric properties and the electrocaloric effect in detail using epitaxial films prepared by pulsed laser deposition (PLD). One of the main challenges is to grow PMN–PT films with a pure perovskite phase as often a secondary phase with a pyrochlore structure is observed, which is regarded as

⁎ Corresponding author at: IFW Dresden, P.O. Box 270116, D-01171 Dresden, Germany. E-mail address: [email protected] (M. Mietschke).

http://dx.doi.org/10.1016/j.tsf.2015.06.032 0040-6090/© 2015 Elsevier B.V. All rights reserved.

the main factor for insufficient dielectric and ferroelectric properties [9, 10]. Our preliminary experiments have shown, that low deposition temperatures and short deposition times, realized by high laser frequencies, are favoring the perovskite phase growth [11]. However, the reproducibility of a stable pure perovskite phase growth is still inadequate, due to a very narrow window of the suitable preparation parameters also reported by other groups [12]. Another option to stabilize the perovskite phase in PMN–PT is the application of miscut substrates [13,14]. It is known from studies on other oxides as CaHfO3, SrRuO3 or BiFeO3, that the local microstructure as well as the functional properties might be tuned by the application of substrates with vicinal surfaces [15,16]. To our knowledge no detailed study was published on the influence of such a vicinal angle on the growth of PMN–PT. Therefore, we will focus in this paper on the effect of different miscut angles of the SrTiO3 (STO) substrate on the structural properties of epitaxial PMN– PT films with a composition close to the morphotropic phase boundary (MPB) (0.31 ≤ x ≤ 0.37). In this case, the coexistence of different structural phases was reported in literature [17]. This leads to the large number of possible directions of the spontaneous polarization and the associated large phase space is expected to increase the size of the electrocaloric effect.

M. Mietschke et al. / Thin Solid Films 589 (2015) 792–797

793

2. Experimental Buffered (001)-oriented STO substrates with miscut angles between 0° and 15° towards the [100] direction were used for the preparation of PMN–PT thin films by PLD using a KrF excimer laser (Lambda Physics) with a wavelength of 248 nm. The films were grown in standard onaxis geometry, where target and substrate surface are parallel. The deposition parameters were selected on the basis of preliminary experiments in order to achieve a pure perovskite phase PMN–PT thin film growth [11]. In a first step, a sintered La0.7Sr0.3CoO3 (LSCO) target was applied to deposit an epitaxial buffer layer on STO, which is used as bottom electrode afterwards. Typically, a temperature of 550 °C, a repetition rate of 10 Hz and an energy density at the target of 1.5 J/cm2 were chosen to grow the LSCO buffer with a thickness of about 20 nm. The PMN–PT was grown with a frequency of 20 Hz and a laser fluence of 1.5 J/cm2 using a commercial single crystalline 0.68Pb(Mg1/3Nb2/3)3– 0.32PbTiO3 target. The PMN–PT film thickness was about 200 nm as determined by a quartz crystal monitor. Furthermore, the film thickness was verified by cross-sectional transmission electron microscopy (TEM) images. The deposition was performed in a vacuum chamber with a base pressure of 10−5 mbar. The buffer layer as well as the PMN–PT film was grown at a constant oxygen pressure of 0.1 mbar. A similar background pressure was used for cooling to room temperature after deposition. The crystalline structure of the PMN–PT films was analyzed with X-ray diffraction (XRD) measurements using Co-Kα radiation in a standard θ–2θ goniometer (Bruker D8 Advance). Additional pole figure measurements were performed with a Philips X'Pert four circle diffractometer using Cu-Kα radiation. The in-plane lattice parameter was determined using reciprocal space maps (RSMs) in a similar diffractometer equipped with a thin film optics. Film lamellas for cross sectional TEM studies were prepared along the miscut direction by the focused ion beam technique out of the central sample area in a FEI Helios 600i. TEM images are taken in a FEI Tecnai T20 microscope equipped with a LaB6 cathode at acceleration voltages of 200 kV. Furthermore, the film composition was examined by using X-ray photoelectron spectroscopy (XPS). The XPS investigations were performed with a PHI 5600 CI system using non-monochromatic Mg-Kα X-rays (350 W) for excitation during depth profiling. Spectra were recorded with a hemispherical electron analyzer working at a pass energy of 29 eV. For depth profiling, Ar+ ions with an energy of 3.5 keV scanned over 2 mm × 2 mm were used. Effective sputtering rates were 0.7 nm/min and 3.7 nm/min measured at a SiO2 reference. The measuring area here was about 800 μm in diameter. For the ferroelectric measurements, 50 nm thick platinum top electrodes were deposited on the film surface at room temperature using a shadow mask with holes of 200 μm in diameter. Temperature and frequency dependent polarization hysteresis loops (P–E) were measured with a Radiant Multiferroic testing system equipped with a Linkam temperature stage in a temperature range between room temperature and 150 °C and in a frequency range between 1 Hz and 100 kHz. 3. Results and discussion 3.1. Structural characterization The XRD patterns of the deposited PMN–PT films grown on LSCO buffered STO substrates are shown in Fig. 1. The miscut of the substrate was corrected by an additional ω-tilt in such way that the (00ℓ) planes of STO are in diffraction condition for these scans. The LSCO buffer layer and the perovskite phase of PMN–PT reveal an (00ℓ) orientation for the different miscut angles on the (001) oriented STO substrate. It is noteworthy that a distinct peak split is visible for the perovskite phase, which will be discussed in detail below. Reflections which correspond to the pyrochlore phase are only visible in the XRD pattern of the film on the standard substrate without a miscut. This indicates that the small deposition window for a pure perovskite growth [12] was not fully met under the used preparation conditions for this sample series.

Fig. 1. Theta–2theta scans of PMN–PT thin films grown on STO/LSCO substrates with different miscut angles. In the measurement geometry the miscut was corrected by an additional ω-tilt.

However, the pyrochlore reflections disappear by using a miscut angle of 5° and higher indicating a substantial increased deposition window for pyrochlore free films. Instead, small additional peaks were observed which can be indexed with all strongest lines of PMN–PT and which can be a hint of a non-epitaxial texture component of the perovskite PMN– PT phase. The intensity of these reflections vanishes with increasing miscut angle and is no longer visible in the 15° miscut film. Pole figure measurements exemplarily shown for a 15° miscut film in Fig. 2a + b were used to verify the epitaxial growth. The epitaxial relationship of the subsequent layers was determined as PMN–PT(001)[100]||LSCO(001)[100]||STO(001)[100] for all samples. The miscut angle of 15° is transferred from the STO substrate to the PMN–PT layer as indicated in the pole figures by a similar shift of the reflections in ψ of about 15°. Subsequent 2 theta/ omega scans for increasing psi angles were performed using an inplane orientation of ϕ = 0° in order to check for a randomly textured PMN–PT component, which would be indistinguishable from the background in pole figure measurements. The results are shown in Fig. 2c + d for the 5° and 15° miscut films, respectively. Besides the expected strong (011) reflections of the STO substrate and the epitaxial PMN–PT layer, an area with increased intensity over the whole psiangle range was found at 2θ = 31.3° for the 5° substrate. This diffraction angle corresponds to the (110) plane of the perovskite structure and points to a randomly textured component in the film already indicated by the additional peaks in the standard XRD scans. A similar line was found in such two-dimensional scans performed at phi-angles, where no epitaxial PMN–PT is visible (i.e., ϕ ≠ 0°, 90°, …). In contrast, no such feature was observed for the 15° miscut film. Cross-sectional TEM images of these two samples shown in Fig. 2e + f reveal a homogeneous film morphology. In both cases a smooth and closed LSCO buffer layer with a thickness of about 20 nm was found. A number of cone shaped precipitates can be observed in the PMN–PT film on the 5° miscut substrate as marked exemplarily in Fig. 2e. Local diffraction investigations in the TEM indicate a different orientation of these grains compared to the epitaxial PMN–PT matrix. Therefore, we assume that these grains correspond to the randomly textured component in this film already detected by XRD. In comparison, none of these grains were observed for the 15° miscut film along the complete lamella (Fig. 2f). In general, the origin of these untextured grains is still unclear and needs a detailed TEM study to unveil the nucleation mechanism for this texture component. The structure of the grown films was characterized in more detail by RSM using the (103) and (002) planes (Fig. 3). The high resolution measurements show clearly three intensity maxima, which are corresponding to STO, LSCO and PMN–PT planes of the film architecture. The LSCO

794

M. Mietschke et al. / Thin Solid Films 589 (2015) 792–797

c)

a)

60

PMN-PT (011) STO (011)

e)

Psi (°)

50 40

Pt

30

PMN-PT LSCO STO

20 10 28

30

32

34

untextured component

2 Theta/omega (°) b)

d)

70

f)

Psi (°)

60 50 40 30 20 28

30

32

34

2 Theta/omega (°) Fig. 2. a) (110) pole figure of a 15° miscut STO substrate; b) (220) pole figure of the deposited PMN–PT film; c) + d) psi-2 theta/omega scans of the PMN–PT film on 5° miscut and 15° miscut STO substrates, respectively; e) + f) TEM cross-section of the PMN–PT film on the 5° and 15° miscut STO, respectively.

peak appears at almost the same Qx value as STO in all films. Since Qx is proportional to the in-plane lattice parameter, it is obvious that LSCO grows coherently strained on the STO substrate similar to previous studies on this material [18,19]. The lattice parameter of LSCO was determined to a = 3.91 Å ± 0.01 Å and c = 3.83 Å ± 0.01 Å, independent of the miscut angle. A small additional tilt of about 0.1° towards the substrate normal was found for the 5° and 10° miscut substrates, which might be explained with the incorporation of dislocations according to the Nagai model to accommodate the lattice misfit [20].

In contrast, the PMN–PT reflections were found at different Qx values compared to the STO substrate, which indicates a relaxed growth. The shape of the PMN–PT reflections is changing with increasing miscut angle and is depending on the measurement direction, i.e., along or perpendicular to the miscut direction. Two clearly separated (103) respectively (002) reflections are visible for the films on the standard as well as on the 5° miscut substrate. However, the distance between the peaks seems to shrink within increasing vicinal angle. The films deposited on higher miscut angle substrates show a broader, unsymmetrical PMN–PT

0.42 )

0.58 LSCO (103)

0° miscut STO (103)

5° miscut

CO LS

2) (0 0

ST O

PMN-PT (002)

0.01 0.02 0.03 Qx (r.l.u.)

-0.01 0 0.01 Qx (r.l.u.) 0.42

d)

e)

0.40

10° miscut PMN-PT (103)

15° miscut

min.

0.46 0.20

0.36

Qz (r.l.u.)

0.50

0.40

0.38

log intensity (a.u.)

a-axis

Qz (r.l.u.)

c-axis

(0

max.

0.62

Qz (r.l.u)

c)

02

a)

0.54

b)

0.25

Qx (r.l.u.)

0.30

0.38

0.36

0.35 -0.01 0 0.01 Qx (r.l.u.)

0.08

0.10 Qx (r.l.u.)

0.12

Fig. 3. a) Reciprocal space maps (RSMs) of the PMN–PT thin films grown on STO/LSCO substrates with different miscut angles; b) + c) RSM of (002) reflections of the 5° miscut film along and perpendicular to the miscut direction, respectively; d) + e) RSM of (002) reflections of the 15° miscut film along and perpendicular to the miscut direction.

M. Mietschke et al. / Thin Solid Films 589 (2015) 792–797

peak, which might be the result of an overlapping of several peaks. Additionally, a significant tilt away from the substrate normal of about 0.4° and 0.8° was found for the layers on the 5° and 10° miscut substrates, respectively (Fig. 3c), whereas no such tilt was observed on the 15° miscut STO. Again, such behavior might be explained by the Nagai model [20] taking into account the larger lattice constant of PMN–PT compared to STO favoring a tilting away from the substrate normal in contrast to the smaller lattice parameter of the LSCO buffer showing an opposite tilt direction. A similar behavior was also observed during the growth of other perovskites as BiFeO3 on vicinal substrates [21]. As the terrace width gets smaller with increasing miscut angles, the higher number of defects might lead to a different behavior for the 15° miscut STO, where no additional tilt was observed. Another possible explanation could be given by incorporation of oxygen vacancies, which leads to a relaxation of the structure whereby no additional tilt is needed [22]. The in-plane and out-of-plane lattice parameters were calculated from the RSM data using the (103) and (103) poles of PMN–PT. The resulting values are summarized in Fig. 4 in dependence of the vicinal angle. No differences in the lattice parameters were found for the measurements along and perpendicular to the miscut direction within the accuracy of our XRD device. It is noticeable that the in-plane lattice parameters (a-axis) are almost constant with increasing vicinal angle and similar to the target value. The films deposited on standard, 5° and 10° miscut STO exhibit a second peak with a significant larger out-of-plane lattice parameter (c-axis), despite an almost similar inplane lattice parameter. The PMN–PT component with a smaller c-axis shows a decreasing out-of-plane lattice parameter with increasing miscut angle resulting in a reduction of the c/a-ratio from 1.02 to 1.002 illustrated in the inset of Fig. 4. Simultaneously, the c/a-ratio changes from 1.05 to about 1.02 for the second PMN–PT component. In comparison to values from the literature [17,23], where a c/a-ratio of 1.012 for a tetragonal structure and of 1.006 for a monoclinic structure is reported, the determined ratio for the first peak in our films might be explained with one of these structural variants. Additionally, the coexistence of both phases for PT contents close to the morphotropic phase boundary needs to be taken into account. Unfortunately, a clear distinction between both phases is not possible due to the resolution limit of the used experimental methods. Furthermore, it is known from literature that the thermal expansion coefficient of cation disordered perovskites as PMN–PT is significantly smaller compared to simple perovskites as STO [24]. This thermal misfit will lead to a higher c/a-ratio after cooling to room temperature due to substrate clamping of

Fig. 4. In-plane and out-of-plane lattice parameters in dependence of the miscut angle of the STO substrate; the inset shows the c/a-ratio for the main peak and the secondary peak of the PMN–PT film in dependence of the miscut angle. The solid line represents the pseudocubic lattice parameter of the target material.

795

the PMN–PT film. However, the discussed effects cannot explain the origin of the second PMN–PT peak on the standard STO substrate with a significant larger c/a-ratio of almost 1.05. The overall stoichiometric composition of the perovskite phase might be different due to the large pyrochlore content in this sample. In general, detailed TEM studies are necessary to clarify the local structure of this film. The composition of the films was studied with XPS measurements in the next step (Fig. 5). For the 0° miscut film, the composition differs from the nominal values of the target compound marked by the dashed lines in the figure in particular for lead and niobium. Such a lead deficiency is often observed in PMN–PT thin films, due to the volatility of lead and lead oxides [2,25]. The lead deficiency is reduced with increasing miscut angle and even a slight lead excess was found. As a result, films grown on the highest miscut angle of 15° exhibit almost the target composition. The results presented so far indicate a dependence of the growth mechanism on the miscut angle of the substrate. A higher miscut is generally connected to smaller terrace widths of the substrate surface, which might favor a step flow growth instead of island growth of the atomic layers [26]. On the other hand, a smaller terrace width also leads to a higher nucleation density as bonding of the deposited atoms is stronger at the step edges. Nevertheless, a change in the growth mechanism might lead to a stabilization of the volatile lead and for this reason a nearly stoichiometric composition of the PMN–PT film on substrates with a higher miscut angle. As a consequence, the formation of the pyrochlore phase, which is lead-deficient, is prevented [13]. 3.2. Ferroelectric characterization Finally, the ferroelectric properties were determined for the grown films. Fig. 6a shows the temperature dependent polarization measurements on the 10° miscut films as one example. The other pure perovskite phase films on the vicinal substrates are showing a similar ferroelectric behavior. At room temperature a remanent polarization of about 2.1 μC/cm2 and a coercive field of 145 kV/cm were determined. The values of the remanent polarization are slightly lower in comparison to published results for thin films with similar PT content, whereas the coercive field is enhanced at room temperature [13,27,28,2,29]. In contrast, pyrochlore-rich films as the grown PMN–PT layer on the standard substrate reveal a high leakage current resulting in a mainly resistive behavior of the films. Therefore, it was not possible to determine reliable polarization loops in this case. This is in good agreement with the literature, where pyrochlore containing PMN–PT films are reported to have very low dielectric constants, mainly due to the high leakage currents in the films [9,10]. The remanent polarization was determined

Fig. 5. XPS measurements of PMN–PT films (sputter depth 70 nm), dashed lines represent the nominal values for target composition of the elements.

796

M. Mietschke et al. / Thin Solid Films 589 (2015) 792–797

Fig. 6. a) Temperature dependent polarization measurements on PMN–PT grown on a 10° miscut substrate, investigated with a maximum electric field of 450 kV/cm and a cycle frequency of 100 Hz; b) remanent polarization as a function of temperature.

as a function of the temperature during heating (Fig. 6b). A sharp drop of the polarization to almost zero was observed at a temperature of approximately 65 °C. This sharp transition from a ferroelectric to paraelectric behavior is characteristic for a first order phase transition, which would point to a normal ferroelectric material in contrast to the typically observed relaxor behavior for this material composition reported by other groups [30]. In comparison to TC values of about 155 °C measured for bulk materials [31], our samples exhibit a lower transition temperature. A similar behavior was already reported by other groups [2]. In this case, different possible origins for the reduced transition temperature were discussed. Whereas effects of the film thickness or the residual stress due to the clamping of the PMN–PT film on the substrate were ruled out by these authors, a dependence of the ferroelectric properties on the crystalline quality was found. This might be also the case for our films as it is known that the deposition conditions during the PLD process can severely influence the local microstructure as the mosaic subgrain structure or the distribution of lattice defects. Furthermore, the oxygen content of the grown film might be sub-optimal after deposition. A subsequent annealing in an oxygen atmosphere would be beneficial in this case to reduce possible oxygen vacancies. Further detailed investigations are necessary to study the influence of these parameters on the ferroelectric properties. Nevertheless, a lower transition temperature might be favorable for the application of this material in electrocaloric solid state cooling devices.

4. Conclusion We have studied the growth of epitaxial PMN–PT thin films on LSCO buffered STO substrates by PLD. PMN–PT films grown on (001) oriented STO substrates exhibit a significant volume fraction of the pyrochlore phase as well as a lead deficiency for the chosen deposition conditions. In contrast, an almost pure perovskite phase with a composition close to the nominal value was observed for PMN–PT layers grown on vicinal STO substrates. This result indicates that the deposition window for pyrochlore-free films is significantly broader, if such substrates are used. In addition, detailed XRD and RSM studies revealed a clear dependence of the structural parameters on the miscut angle implying a substantial influence of the substrate surface structure on the nucleation and growth mechanism of PMN–PT. Whereas the majority of the measured XRD peaks might be explained by the coexistence of monoclinic and tetragonal phases for compositions close to the morphotropic phase boundary, a secondary structure with a high c/a-ratio was observed for films on standard substrates. Further detailed high resolution

TEM studies are essential to clarify the microstructure of the PMN–PT films on the local scale and to verify the existence of different structural phases. Nevertheless, the variation of the vicinal angle opens also the opportunity to tune the crystalline structure of PMN–PT for optimized functionalities. Finally, all pure perovskite phase PMN–PT films on vicinal STO show a comparable ferroelectric behavior with PR of about 2 μC/cm2, EC of 145 kV/cm and TC of about 65 °C. Acknowledgments The work is partially funded by the DFG under grant no. HU1726/3 in the framework of the priority program SPP 1599 Ferroic cooling. The authors thank U. Besold for the technical support as well as A. Pöhl and T. Sturm for their preparation of the FIB lamellas. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]

[12] [13] [14]

[15] [16] [17] [18] [19] [20] [21] [22]

J.F. Scott, Annu. Rev. Mater. Res. 41 (2011) 229–240. J.-P. Maria, W. Hackenberger, S. Trolier-McKinstry, J. Appl. Phys. 84 (1998) 5147–5154. S.-E. Park, T.R. Shrout, J. Appl. Phys. 82 (1997) 1804. J. Wang, K. Wong, H. Chan, C. Choy, Appl. Phys. A 79 (2004) 551–556. S. Zhang, F. Li, J. Appl. Phys. 111 (2012) 031301. S. Fähler, U.K. Rößler, O. Kastner, J. Eckert, G. Eggeler, H. Emmerich, P. Entel, S. Müller, E. Quandt, K. Albe, Adv. Eng. Mater. 14 (2012) 10–19. T.M. Correia, J.S. Young, R.W. Whatmore, J.F. Scott, N.D. Mathur, Q. Zhang, Appl. Phys. Lett. 95 (2009) 182904. J. Hagberg, A. Uusimaki, H. Jantunen, Appl. Phys. Lett. 92 (2008) 132909. J.P. Guha, D.J. Hong, H.U. Anderson, J. Am. Ceram. Soc. 71 (1988) C-152–C-154. J. Chen, M.P. Harmer, J. Am. Ceram. Soc. 73 (1990) 68–73. M. Mietschke, S. Fähler, L. Schultz, R. Hühne, in: Proc. IEEE International Symposium on Application of Ferroelectrics - InternationalWorkshop on Acoustic Transduction Materials and Devices —, Piezoforce Microscopy (2014) pp. 153–156, http://dx.doi.org/10.1109/ ISAF.2014.6922996. M. Boota, E.P. Houwman, M. Dekkers, M. Nguyen, G. Rijnders, Appl. Phys. Lett. 104 (2014) 182909. S.D. Bu, M.K. Lee, C.B. Eom, W. Tian, X.Q. Pan, S.K. Streiffer, J.J. Krajewski, Appl. Phys. Lett. 79 (2001) 3482–3484. S.H. Baek, J. Park, D.M. Kim, V.A. Aksyuk, R.R. Das, S.D. Bu, D.A. Felker, J. Lettieri, V. Vaithyanathan, S.S.N. Bharadwaja, N. Bassiri-Gharb, Y.B. Chen, H.P. Sun, C.M. Folkman, H.W. Jang, D.J. Kreft, S.K. Streiffer, R. Ramesh, X.Q. Pan, S. Trolier-McKinstry, D.G. Schlom, M.S. Rzchowski, R.H. Blick, C.B. Eom, Science 334 (2011) 958–961. B.W. Lee, C.U. Jung, M. Kawasaki, Y. Tokura, J. Appl. Phys. 104 (2008) 103909. H.W. Jang, D. Ortiz, S.-H. Baek, C.M. Folkman, R.R. Das, P. Shafer, Y. Chen, C.T. Nelson, X. Pan, R. Ramesh, C.-B. Eom, Adv. Mater. 21 (2009) 817–823. B. Noheda, D. Cox, G. Shirane, J. Gao, Z.-G. Ye, Phys. Rev. B: Condens. Matter Mater. Phys. 66 (2002) 1–10 054104. A.D. Rata, A. Herklotz, K. Nenkov, L. Schultz, K. Dörr, Phys. Rev. Lett. 100 (2008) 076401. S. Trommler, R. Hühne, A. Rata, L. Schultz, B. Holzapfel, Thin Solid Films 519 (2010) 69–73. H. Nagai, J. Appl. Phys. 45 (1974) 3789–3794. T.H. Kim, S.H. Baek, S.Y. Jang, S.M. Yang, S.H. Chang, T.K. Song, J.-G. Yoon, C.B. Eom, J.-S. Chung, T.W. Noh, Appl. Phys. Lett. 98 (2011) 022904. S. Pennycook, H. Zhou, M. Chisholm, A. Borisevich, M. Varela, J. Gazquez, T. Pennycook, J. Narayan, Acta Mater. 61 (2013) 2725–2733.

M. Mietschke et al. / Thin Solid Films 589 (2015) 792–797 [23] Y.M. Jin, Y.U. Wang, A.G. Khachaturyan, J.F. Li, D. Viehland, Phys. Rev. Lett. 91 (2003) 197601. [24] K. Uchino, S. Nomura, A. Amin, Z.P. Chang, L.E. Cross, R.E. Newnham, Jpn. J. Appl. Phys. 19 (1980) L398–L400. [25] Z. Surowiak, M. Kupriyanov, A. Panich, R. Skulski, J. Eur. Ceram. Soc. 21 (2001) 2783–2786. [26] Q. Gan, R.A. Rao, C.B. Eom, Appl. Phys. Lett. 70 (1997) 1962–1964.

797

[27] Z. Feng, D. Shi, R. Zeng, S. Dou, Thin Solid Films 519 (2011) 5433–5436. [28] J. Jiang, H.-H. Hwang, W.-J. Lee, S.-G. Yoon, Sensors Actuators B Chem. 155 (2011) 854–858. [29] S.A. Yang, S.Y. Cho, J.S. Lim, S.D. Bu, Thin Solid Films 520 (2012) 7071–7075. [30] G.A. Samara, J. Phys. Condens. Matter 15 (2003) R367. [31] S.W. Choi, T.R. Shrout, S.J. Jang, A.S. Bhalla, Ferroelectrics 100 (1989) 29–38.