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3)O3-PbTiO3 thin films
Materials Characterization 129 (2017) 234–241
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Effect of substrate miscut on the microstructure in epitaxial Pb(Mg1/3Nb2/3) O3-PbTiO3 thin films
MARK
Paul Chekhonina,c,⁎, Michael Mietschkea,b, Darius Pohla, Frank Schmidta,b, Sebastian Fählera, Werner Skrotzkic, Kornelius Nielscha,b, Ruben Hühnea a b c
Institute for Metallic Materials, IFW Dresden, P.O. Box 270116, D-01171 Dresden, Germany Institut für Werkstoffwissenschaft, Technische Universität Dresden, Helmholtzstraße 7, D-01069 Dresden, Germany Institut für Strukturphysik, Technische Universität Dresden, Haeckelstraße 3, D-01062 Dresden, Germany
A R T I C L E I N F O
A B S T R A C T
Keywords: PMN-PT Ferroelectric Thin films Microstructure Dislocation structure Stacking faults
Pb(Mg1/3Nb2/3)O3-PbTiO3 is known for its excellent ferroelectric and electrocaloric properties. Epitaxial films, which allow applying high voltages, can be grown in a much broader parameter range when depositing them on miscut substrates than on those without miscut. To understand the influence of substrate miscut, here the structure and microstructure of two films, with and without miscut, is compared in detail. In both cases the epitaxial strain is relaxed, but while in films without miscut perfect misfit dislocations are observed, the film with miscut exhibits many partial dislocations and stacking faults. The stacking faults may play an important role for the stabilization of the pure perovskite Pb(Mg1/3Nb2/3)O3-PbTiO3 phase in thin films on miscut substrates.
1. Introduction
LaAlO3). Additionally, partial dislocations with Burgers vectors b = a / 2〈110〉 resulting in a stacking fault (SF) in-between are observed in some of these materials [12,13,14,15]. Such O-sublattice preserving stacking faults have also been reported for bulk 0.65PMN-0.35PT [16]. Whereas a significant number of microstructural data is available for simple perovskites, such detailed knowledge is lacking for PMN-PT thin films. Typically, integral X-ray diffraction measurements are used to provide evidence for the growth of a single phase material and to demonstrate the desired epitaxial relationship, whereas transmission electron microscopy (TEM) cross sections give an overview over the general microstructure. However, to our best knowledge no detailed microstructural investigations focusing on lattice defects in PMN-PT thin films were published so far. Just in reference [17] misfit dislocations are reported for 0.67PMN-0.33PT thin films grown on La0.5Sr0.5CoO3/CeO2/YSZ buffered silicon without giving further details. Therefore, our aim is to perform such detailed microstructural investigations on epitaxial 0.68PMN-0.32PT thin films grown on La0.7Sr0.3CoO3 (LSCO) buffered SrTiO3 (STO). In general, the parameter window for the growth of pure perovskite PMN-PT thin films without the lead-deficient pyrochlore phase is quite narrow, if pulsed laser deposition (PLD) is applied [21,22]. However, it was demonstrated that the growth on STO substrates with a defined
The perovskite (1 - x)Pb(Mg1/3Nb2/3)O3-xPbTiO3 (PMN-PT) shows appealing electrocaloric properties [1,2] in addition to the well-known large dielectric constants and high piezoelectric coefficients [3,4] and therefore may become a suitable material for cooling applications [5]. So far, many studies were conducted on bulk PMN-PT [6,7] as well as on thin films [8,9,10] to determine the caloric properties. Among them, thin films of PMN-PT are of great interest because of the high applicable electric fields [6] and the possibility to modify the crystal orientation and strain state through epitaxy. At the same time, a detailed microstructural analysis is required to understand the influence of the growth parameters and the resulting structure on the functional properties of the films. Therefore, we focus in this report on the defect structures of epitaxial PMN-PT films grown on SrTiO3 substrates with different miscut. Microstructural investigations on lattice defects of epitaxial ABO3 perovskites thin films typically show that these materials contain perfect misfit dislocations with Burgers vectors b = a〈100〉 to accommodate the elastic strain (for example for SrZrO3 [11], BaZrO3 [11], SrTi0.5Zr0.5O3 [11], BaTiO3 [12] and PbSc0.5Ta0.5O3 [13] grown on SrTiO3 as well as for Ba0.75Sr0.25TiO3 [14] and SrTiO3 [15] grown on
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Corresponding author at: Institute for Metallic Materials, IFW Dresden, P.O. Box 270116, D-01171 Dresden, Germany. E-mail addresses: [email protected] (P. Chekhonin), [email protected] (M. Mietschke), [email protected] (D. Pohl), [email protected] (F. Schmidt), [email protected] (S. Fähler), [email protected] (W. Skrotzki), [email protected] (K. Nielsch), [email protected] (R. Hühne). http://dx.doi.org/10.1016/j.matchar.2017.05.003 Received 21 December 2016; Received in revised form 28 March 2017; Accepted 6 May 2017 Available online 07 May 2017 1044-5803/ © 2017 Elsevier Inc. All rights reserved.
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Fig. 1. XRD of PMN-PT thin films grown on LSCO buffered STO substrates: a) Θ-2Θ-scans (A10 was mounted with an inclination to compensate the 10° substrate tilt) and b) to e) pole figures of both samples. b) and d) show the (110) pole figures of STO of samples A0 and A10, respectively, while c) and e) depict (220) pole figures of PMN-PT of A0 and A10, respectively.
2.2. Instrumentation
miscut is helpful to stabilize the desired perovskite PMN-PT phase [18,19]. In this work, two pure perovskite PMN-PT thin film samples are compared, which were prepared on a STO substrate without as well as with a 10° miscut using slightly different deposition parameters. The central aim of this study is to depict microstructural differences between these two samples in order to clarify the influence of the miscut on the film growth. In the following, the perovskite PMN-PT phase is referred to just as PMN-PT. Additionally, it should be noted that the crystallographic structure of bulk 0.68PMN-0.32PT is not cubic at room temperature [20]. This composition is close to the so-called morphotropic phase boundary, where monoclinic, rhombohedral or tetragonal lattice structures coexist. However, the quantitative differences between these structures are small, so that PMN-PT will be treated as a cubic lattice with respect to crystallographic directions and planes.
The thin film samples were analyzed by X-ray diffraction (XRD). Θ2Θ-scans were performed with a Bruker D8 Advance diffractometer and pole figures were recorded in a Philips X'Pert four circle diffractometer using Co-Kα and Cu-Kα radiation, respectively. The determination of inplane lattice parameters was done by recording reciprocal space maps (RSMs) in an additional Philips X'Pert diffractometer equipped with a thin film optics using a Goebel mirror in the primary beam and narrow Soller slits on the detector side. Surface characterization was carried out with a Bruker Dimension Icon atomic force microscope (AFM). Transmission electron microscopy investigations were done in a FEI Tecnai T20 TEM equipped with a LaB6 cathode operated at 200 kV acceleration voltage and an aberration corrected FEI Titan3 80–300 TEM operated in scanning TEM (STEM) mode at 300 kV using a high angle annular dark-field (HAADF) detector. Cross-sectional TEM lamellae were prepared out of the central areas of both samples applying the focussed ion beam (FIB) technique in a FEI Helios 600i FIB and subsequent argon ion milling. The image plane in all TEM lamellae is (010). The Burgers vectors for dislocations are typically determined by a Burgers circuit in the HAADF images.
2. Experimental 2.1. Sample Preparation 0.68PMN-0.32PT thin films were deposited on single crystalline STO substrates applying pulsed laser deposition using a 30 nm to 50 nm thin LSCO buffer layer. One sample (referred to as A0) was produced on a STO substrate with the normal direction parallel to [001]. For A0 the epitaxial relationship along the film plane is: STO (001) || LSCO (001) || PMN-PT (001) and along an in-plane direction: STO [100] || LSCO [100] || PMN-PT [100]. Another sample (referred to as A10) was grown on a (001)-oriented STO substrate with a 10° miscut towards the [100] direction showing the same epitaxial relationship. The lattice parameters of bulk STO, LSCO and PMN-PT are 3.905 Å, 3.839 Å and about 4.02 Å, respectively. Therefore, in the case of epitaxy without strain relaxation, a biaxial in-plane tensile strain and biaxial in-plane compressive strain is introduced to the LSCO and PMN-PT layers, respectively. The deposition of both samples was performed in an in-house assembled PLD chamber equipped with a KrF excimer laser (248 nm wavelength, Lambda Physics) at a temperature of 550 °C. Sintered La0.7Sr0.3CoO3 and single crystalline 0.68PMN-0.32PT disks were used as targets. The deposition of sample A10 was performed at constant oxygen pressure of 0.1 mbar and a laser energy density on the target surface of about 1.5 J/cm2, more details are published in reference [19]. In contrast, sample A0 was prepared with an energy density of 0.8 J/cm2 and an oxygen pressure of 0.3 mbar to meet the narrow deposition window for standard substrates. Further details can be found in [22]. After deposition, both samples were transferred to a furnace and annealed for 1 h at 400 °C in pure oxygen at atmospheric pressure to reduce oxygen vacancies.
3. Results 3.1. X-ray Diffraction XRD Θ-2Θ-scans reveal the successful growth of LSCO and the perovskite PMN-PT phase (Fig. 1 a)) in both samples. Essentially, the (00ℓ) peaks of PMN-PT, STO and LSCO are clearly visible, whereas no sign for a pyrochlore structure was found. Some traces of non-epitaxial PMN-PT are observable as marked in Fig. 1 a). This is in accordance with the presence of a small number of protruding grains at the PMN-PT film surface (cf. one example in Fig. 4 a)) with different orientation. The AFM maps in Fig. 2 show that such grains exist in both samples, but apparently have a much higher area density in A10 in comparison to A0. However, the volume density is in general quite low. (110) pole figures of STO and (220) pole figures of PMN-PT prove the above mentioned epitaxial relationship for sample A0 (Fig. 1 b) and c)) and A10 (Fig. 1 d) and e)), respectively. Measured (110) pole figures of the LSCO layers (not shown here) are qualitatively similar to the (110) STO pole figures of the corresponding sample. RSMs of (103) reflections provide more details on the structural properties of the layers (Fig. 3). In each RSM, three distinct intensity maxima are visible, which arise from the STO, LSCO and PMN-PT layer. The in-plane lattice parameter (a-axis) of the layers, which correspond to the Qx values of the reciprocal lattice, were calculated using the (103) and (103) maxima, while the out of plane lattice parameter (caxis) was extracted from corresponding (002) peaks of the Θ-2Θ-scans. The results are listed in Table 1. 235
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Fig. 2. AFM mappings of a) A0 and b) A10. The bright spots are protruding grains.
Within the measurement uncertainty of RSM, the in-plane lattice parameter of LSCO is similar in A0 and A10 and equal to the value of STO (a = 3.905 Å), thus indicating zero or negligible strain relaxation. The c-axis of the LSCO layers shrinks as expected in the presence of biaxial in-plane tensile strain. On the other hand, the larger in-plane lattice parameter of PMN-PT indicates nearly complete strain relaxation. The slightly larger PMN-PT c-axis for both samples may be explained either by some remaining compressive in-plane strain that may be affected by oxygen vacancies or by a non-cubic crystal structure [20]. Obviously, the in-plane strain in PMN-PT is not affected by the substrate miscut. It was concluded from the RSM measurements that the [001] directions of the substrate and each subsequent layer are aligned exactly parallel in sample A0. This is not the case in A10, where the LSCO layer is slightly tilted towards the substrate normal by about 0.1° (positive tilt), while the PMN-PT layer is tilted by approximately −0.8° in the opposite direction (i.e. away from the substrate normal). These tilts might be explained qualitatively as a consequence of a stepped surface by the Nagai model [23]. In this case, the tilt angle α can be calculated by the Nagai Eq. (1)
⎛ c − cB ⎞ tan α = −⎜ E ⎟ tan φ ⎝ cB ⎠
Table 1 Lattice constants (a and c) and film thickness (t) determined by RSM, XRD and TEM, respectively, of the LSCO and PMN-PT layers. A0 aLSCO in Å (by RSM) cLSCO in Å (by Θ-2Θ) aPMN-PT in Å (by RSM) cPMN-PT in Å (by Θ-2Θ) tLSCO in nm (by TEM) tPMN-PT in nm (by TEM)
3.91 ± 3.78 ± 4.02 ± 4.03 ± ≈55 ≈160
A10 0.01 0.01 0.02 0.01
3.91 ± 3.83 ± 4.02 ± 4.06 ± ≈ 30 ≈ 300
0.01 0.01 0.02 0.01
with the measured lattice constants for sample A10 from Table 1. It is noticeable that the calculated tilt of the PMN-PT layer is lower than the measured one, which might be explained by the observed microstructure described later. It should be mentioned in this context that even an extended Nagai model, which is based on a step configuration with different step heights [24], fails to explain the higher tilt measured for sample A10.
3.2. Microstructure (1)
Cross sectional TEM bright field images with the electron beam parallel to the [010] zone axis of PMN-PT are shown in Fig. 4. The images reveal clear interfaces between the distinct layers for samples A10 (Fig. 4 a)) and A0 (Fig. 4 e) ), respectively. The layer thickness of LSCO and PMN-PT shows only minor variations in the nm range. The LSCO layer for A0 is with about 55 nm considerably thicker than
where φ is the miscut angle and cE and cB are the out-of-plane lattice parameters of the tilted epitaxial layer and the substrate or buffer layer beneath, respectively. The theoretical tilt angle towards the STO substrate normal is 0.2° for LSCO and − 0.4° (=−0.6° + 0.2° (considering the tilt of the LSCO layer beneath)) for PMN-PT using Eq. (1)
Fig. 3. RSMs of a) sample A0 and b) A10 showing the (103) reflexions of PMN-PT, LSCO and STO.
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Fig. 4. a) to c) show bright field TEM images of sample A10 at different magnifications highlighting the lamellar structure. In d) a HAADF STEM image demonstrates a stacking fault (SF), its projection is marked by a red line. e) and f) show bright field TEM images of sample A0. The [010] zone axis is always parallel to the electron beam. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
10 of these with b = a / 2[101] (Fig. 5 b)), 4 with b = a / 2[101] (Fig. 5 c)) and 4 with b = a[100]. Fig. 5 d) shows the latter case in the vicinity of a SF, which starts already in the LSCO layer. Despite the small number of identified cases, one may presume that the partial MDs at the interface are obviously unequally distributed, with b = a / 2[101] preferred to b = a / 2[101]. This is comprehensible, as the sample symmetry for a (001) plane is not granted due to the miscut. As a consequence of partial MDs at the LSCO/PMN-PT interface, SFs are introduced to the LSCO layer of sample A10. Representative HAADF images of the STO/LSCO interface shown in Fig. 6 indicate an ideal coherent growth for both samples. While the brighter La/Sr columns are clearly distinguished from the darker Co columns in the complete LSCO layer of A0, this is only the case directly at the STO/LSCO interface for sample A10. Lattice distortions are observed in the LSCO layer just a few nm above that interface. As a result, the contrast of the columns changes continuously along a chosen atomic plane. As stated above, SFs are apparently present inside the LSCO layer of sample A10, however, (partial) dislocations could not be revealed. Therefore, the SFs in the LSCO layer of sample A10 are presumably continuous and bordered by partial MDs in PMN-PT at the LSCO/PMN-PT interface. Whereas extended PMN-PT SFs start occasionally already in the LSCO layer (Fig. 5 d)), most of them appear above the LSCO/PMN-PT interface of sample A10 bordered by a partial dislocation at this position (Fig. 7 a) and b)). Five cases were identified: In three cases the Burgers vector of the partial dislocations is b = a / 2[101] and in the two other cases b = a / 2[101] and b = a / 2[101]. Some of the SFs do not extend up to the sample surface, but are bordered by another partial dislocation, two examples are shown in Fig. 7 c) and d). Here also five cases were identified, four of them with b = a / 2[101] and one with b = a / 2[101]. Finally, perfect dislocations with b = a[001] are also present in the PMN-PT layer of A10 (Fig. 7 c)). In contrast to A10, no partial MDs were observed at the LSCO/PMNPT interface for sample A0. Instead, partial dislocations very rarely interrupt the otherwise widely extended zig-zag shaped SFs in the PMNPT layer of this sample. The lattice appears to be noteworthy bend in the vicinity of SFs in
intended resulting in some cracks. However, the RSM results indicate that these cracks do not lead to a significant strain relaxation. Lamellar aligned structures are revealed in the PMN-PT layer of sample A10 (Fig. 4 b)), whereupon HAADF STEM images (Fig. 4 d)) indicate that they originate from stacking faults. Since the HAADF intensity strongly depends on the atomic number, Pb columns appear as bright spots, while columns with Mg, Nb and Ti atoms appear less bright due to the lower scattering rate resulting from their lower density. Therefore, SFs can easily be recognized by a lattice displacement along the a / 2〈101〉 direction. This is especially the case in few locations, where the normal direction of the SF is perpendicular to [010], as exemplary marked by a red line in Fig. 4 d). However, the normal direction of the SF is typically not aligned perpendicular to [010]; therefore, they appear broadened in the TEM images. The majority of the SFs in A10 have a preferred alignment. The typical angle between the SF normal direction (more accurately its projection in the (010) plane) and the thin film normal is about 40° and therefore about 30° with respect to [001], considering the 10° substrate miscut in this sample. Their alignment is never straight but seems to change direction. Also the width of the projection is different at different locations. Thus, a preferred crystallographic alignment seems to be absent. The distance between two SFs is of the order of 15 nm along the thin film normal, variations of a few nm are typical. As can be seen in Fig. 4 c), the SFs do not start immediately at the LSCO/PMN-PT interface. SFs are also present in sample A0, but here they appear in a zig-zag-arrangement (Fig. 4 f)) at the centre or the upper part of the PMN-PT layer. Throughout the thickness of the PMN-PT layer, never more than one SF is present at a certain position. Therefore, the density of SFs in A0 is about one order of magnitude lower compared to sample A10. A closer look at the LSCO/PMN-PT interface of A0 reveals that strain relaxation is realized only by misfit dislocations (MDs) with Burgers vectors b = a[100]. A typical example is shown in Fig. 5 a). In contrast, such perfect in-plane MDs are not the rule at the LSCO/PMN-PT interface of A10. Instead, the interface is dominated by partial MDs, where Fig. 5 b) is presenting a typical case. Along the LSCO/PMN-PT interface of A10 18 MDs were identified, 237
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Fig. 5. HAADF STEM images of interface dislocations at the LSCO/PMN-PT boundary in a) A0 and b) to d) A10.
(for example the Pb columns below the red line, marked by yellow arrows in Fig. 8), so that exactly at the SF the displacement is a / 2[110]. It should be noted that the STEM micrographs illustrated in this work are excerpts of large scale, high resolution HAADF images having larger regions with blurred appearance due to strain. Therefore, an exact dislocation density cannot be provided with confident statistical
many locations of both samples (Fig. 8). In the case of a SF with a lattice displacement of a / 2[101], Pb columns are expected to shift by a / 2[100] and a / 2[001]. However, if a larger lattice area is observed along a straight line, the lattice translation is not always a / 2[001] perpendicular to that line, but has a smaller value or is even zero as demonstrated by the red line crossing a SF in Fig. 8. This is not obvious, when following the lattice from one column to the next across the SF
Fig. 6. HAADF STEM images of the STO/LSCO interface of a) A10 an b) A0.
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Fig. 7. HAADF STEM images of partial dislocations in the PMN-PT layer of A10 bordering SFs.
4. Discussion If a film exceeds a certain thickness, strain relaxation may occur by MDs [25]. This is apparently the case for both PMN-PT layers investigated. In ref. [16] it was reported that (partial) dislocations in PMN-PT remain unaffected by the electron beam induced heating. Therefore it is concluded that the formation of dislocations and SFs must have taken place during film growth or during annealing at elevated temperatures. As the dislocation energy scales with b2, dissociation of perfect dislocations with Burgers vector b = a〈100〉 into two partial dislocations with b = a / 2〈110〉 is energetically equivalent. However, additional energy is stored in the SF bordered by the partial dislocations. The frequent occurrence of SFs in A10 and their irregular alignment qualitatively indicates that their specific energy seems to be low and not strongly anisotropic. The predominant presence of partial dislocations at the LSCO/PMNPT interface of A10 might be caused by a different mechanism of MD formation as for sample A0, thus resulting in the observed microstructural differences. One argument for such a difference includes the miscut in A10 that reduces the sample symmetry. Another argument considers the smaller lattice constant of LSCO and the atomic steps in A10 at its LSCO/PMN-PT interface, which lead to an additional compressive strain along [001] in the PMN-PT layer. This strain may trigger a relaxation mechanism, which preferably creates partial MDs with b = a / 2[101]. The latter argument does consider that the mechanism responsible for the formation of partial MDs in A10 cannot only be the dissociation of perfect dislocations with b = a[100] or
Fig. 8. HAADF STEM images of a SF in A10. The red dotted line demonstrates the lattice bending close to the SF. Close to the SF yellow arrows mark Pb columns that indicate a lattice displacement of a / 2[110]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
reliability without significant additional experimental effort, which was beyond the scope of our studies. However, SFs with a lattice displacement vector of a / 2〈100〉 or Burgers circuits with b = a / 2〈100〉 were not observed.
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strained thin films [31].
b = a[101], as this would lead to a balanced distribution between partial MDs with b = a / 2[101] and b = a / 2[101] at the LSCO/PMNPT interface. To observed differences between samples A0 and A10 concerning the misfit dislocation structure are not clarified in this work and require an advanced analysis focusing on the dislocation nucleation mechanism. From the experimental data gathered for A10 it can be concluded that partial dislocations bordering the extended SFs in PMN-PT do not contribute to a unilateral strain relaxation or any lattice rotation, since the Burgers vector sum of the 10 partial dislocations observed is equal to zero. However, this is not the case for the MDs at the LSCO/PMN-PT interface, where the excess of b = a / 2[101] MDs should result in a positive lattice rotation about [010]. Even though the experimental data do not provide sufficient statistical reliability, this rotation is probably overcompensated by the existence of perfect dislocations with b = a[001], so that the overall negative rotation observed by XRD can be understood. The 10 partial dislocations enclosing the SFs in A10 indicate that strain accommodation is not responsible for their existence. In general, SFs with a displacement vector of a / 2〈110〉 give rise to nonstoichiometry, if the normal direction of a SF with a displacement vector of a / 2〈110〉 in ABO3 perovskite is not aligned with 〈110〉 [26]. In this case, a SF can accommodate BO2 excess or AO deficit in ABO3. Such SFs have been observed in non-stoichiometric TiO2-rich homoepitaxial SrTiO3 thin films [27]. This leads to a theory, why the undesired lead deficient pyrochlore phase is hampered on miscut substrates: The formation of SFs is preferred due to the miscut. If the deposition parameters result in some lead deficiency, this deficiency may be accommodated by a dense network of extended SFs, thus preventing the formation of the lead deficient pyrochlore phase. Indeed, the deposition parameters for sample A0 had to be carefully adjusted to avoid the pyrochlore phase. A sample deposited on a substrate without miscut exhibits the undesired pyrochlore phase, if produced with the same deposition parameters like sample A10 [19]. This argument is supported by the fact that the observed SFs in A10 are not aligned with 〈110〉 (hence non-stoichiometric) and that the observed partial dislocations enclosing the SFs in A10 are collectively strain neutral. The preferred alignment of the SFs in A10 indicates that their formation is obviously related to the sample symmetry. Some of the SFs may be initiated by partial misfit dislocations that exist in this sample, but most of the partial dislocations bordering a SF are clearly no misfit dislocations (e.g. Fig. 7 b)). Therefore, a different mechanism must be responsible for their formation, which cannot be provided here yet. An advanced microstructural analysis of PMN-PT films with very small miscut angles or lead rich targets might be helpful to clarify this issue. It should be noted in this context that an integral X-ray photoelectron spectroscopy measurement of sample A10 published in ref. [19] did not indicate a larger lead deficit. However, smaller local deviations cannot be excluded. This work does not include measurements regarding ferroelectric and electrocaloric properties. It is worth mentioning that SFs observed in ordered bulk Pb(Sc1/2Ta1/2)O3 seem not to interact with ferroelectric domains [28]. This could also be the case here, however it may be assumed that in the presented PMN-PT thin film A10 the high density of SFs may affect the mobility of ferroelectric domain boundaries. Advanced research is required to analyze this issue in PMN-PT thin films. In both samples it was impossible to determine the precise lattice symmetry, which presumably is not cubic [20]. The mosaicity of the PMN-PT lattice made TEM diffraction experiments difficult and the SFs (especially in A10) lead to broad peak splitting in the TEM diffraction plane. Useful (large angle) convergent beam electron diffraction patterns, which might be used for lattice symmetry determination in PMN-PT (following the work of Wang et al. [29,30] in bulk PMN-PT), were impossible to record. This is mainly due to bending of the TEM lamellae caused by strain relaxation, which is known to occur in
5. Conclusions Two 0.68PMN-0.32PT thin films were produced on La0.7Sr0.3CoO3 (LSCO) buffered (001) oriented SrTiO3 substrates using pulsed laser deposition, one of them with a 10° miscut towards the [100] direction. It is shown that minor changes in substrate orientation can lead to significant differences in the microstructure. In the sample without substrate miscut, misfit dislocations have a Burgers vector b = a[100], while SFs and partial dislocations are absent in the vicinity of the LSCO/PMN-PT interface. In the centre or the upper part of its PMN-PT layer SFs are sporadically present and only a few partial dislocations were observable. In strong contrast, in the sample grown on a miscut substrate the majority of misfit dislocations are partial dislocations with b = a / 2[101] or b = a / 2[101], while there is a strain induced surplus of the former case. The microstructure of its PMN-PT layer is dominated by closely aligned SFs that are bordered by partial dislocations. Presumably, the miscut results in a favoured formation of extended non-stoichiometric SFs, whose formation is not yet fully understood. Furthermore, the authors suggest that these SFs may play an important role in the accommodation of a lead deficiency, which might originate from the preparation of PMN-PT films using non-optimized deposition parameters. Consequently, the formation of the undesired lead deficient pyrochlore phase is impeded and the perovskite PMN-PT phase is stabilized in thin films on miscut substrates. Further research is suggested to clarify this issue. Acknowledgements 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 technical support as well as A. Pöhl for preparation of the TEM lamellae. References [1] J.F. Scott, Electrocaloric materials, Annu. Rev. Mater. Res. 41 (2011) 229–240. [2] X. Moya, S. Kar-Narayan, N.D. Mathur, Caloric materials near ferroic phase transitions, Nat. Mater. 13 (2014) 439–450. [3] S.-E. Park, T.R. Shrout, Ultrahigh strain and piezoelectric behavior in relaxor based ferroelectric single crystals, J. Appl. Phys. 82 (1997) 1804–1811. [4] J. Wang, K.H. Wong, H.L.W. Chan, C.L. Choy, Composition control and electrical properties of PMN-PT thin films around the morphotropic boundary, Appl. Phys. A Mater. Sci. Process. 79 (2004) 551–556. [5] S. Fähler, U.K. Rößler, O. Kastner, J. Eckert, G. Eggeler, H. Emmerich, P. Entel, S. Müller, E. Quandt, K. Albe, Caloric effects in ferroic materials: new concepts for cooling, Adv. Eng. Mater. 14 (2012) 10–19. [6] J. Hagberg, A. Uusimäki, H. Jantunen, Electrocaloric characteristics in reactive sintered 0.87 Pb(Mg1 ∕ 3Nb2 ∕ 3)O3–0.13 PbTiO3, Appl. Phys. Lett. 92 (2008) 132909. [7] B. Rožič, M. Kosec, H. Uršič, J. Holc, B. Malič, Q.M. Zhang, R. Blinc, R. Pirc, Z. Kutnjak, Influence of the critical point on the electrocaloric response of relaxor ferroelectrics, J. Appl. Phys. 110 (2011) 064118. [8] A.S. Mischenko, Q. Zhang, R.W. Whatmore, J.F. Scott, N.D. Mathur, Giant electrocaloric effect in the thin film relaxor ferroelectric 0.9 PbMg1/3Nb2/3O3–0.1 PbTiO3 near room temperature, Appl. Phys. Lett. 89 (2006) 242912. [9] T.M. Correia, J.S. Young, R.W. Whatmore, J.F. Scott, N.D. Mathur, Q. Zhang, Investigation of the electrocaloric effect in a PbMg2/3Nb1/3O3-PbTiO3 relaxor thin film, Appl. Phys. Lett. 95 (2009) 182904. [10] D. Saranya, A.R. Chaudhuri, J. Parui, S.B. Krupanidhi, Electrocaloric effect of PMN–PT thin films near morphotropic phase boundary, Bull. Mater. Sci. 32 (2009) 259–262. [11] P.A. Langjahr, F.F. Lange, T. Wagner, M. Rühle, Lattice mismatch accommodation in perovskite films on perovskite substrates, Acta Mater. 46 (1998) 773–785. [12] T. Suzuki, Y. Nishi, M. Fujimoto, Analysis of misfit relaxation in heteroepitaxial BaTiO3 thin films, Phil. Mag. A 79 (1999) 2461–2483. [13] B.I. Birajdar, A. Chopra, M. Alexe, D. Hesse, Crystal defects and cation ordering domains in epitaxial PbSc0.5Ta0.5O3 relaxor ferroelectric thin films investigated by high-resolution transmission electron microscopy, Acta Mater. 59 (2011) 4030–4042. [14] Y.Q. Wang, W.S. Liang, P.K. Petrov, N.M. Alford, Dissociation of misfit and threading dislocations in Ba0.75Sr0.25TiO3 epitaxial film, Mater. Charact. 62 (2011) 294–297. [15] Y.L. Qin, C.L. Jia, K. Urban, J.H. Hao, X.X. Xi, Dislocations in SrTiO3 thin films
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Effect of substrate miscut on the microstructure in epitaxial Pb(Mg1/3Nb2/3) O3-PbTiO3 thin films
MARK
Paul Chekhonina,c,⁎, Michael Mietschkea,b, Darius Pohla, Frank Schmidta,b, Sebastian Fählera, Werner Skrotzkic, Kornelius Nielscha,b, Ruben Hühnea a b c
Institute for Metallic Materials, IFW Dresden, P.O. Box 270116, D-01171 Dresden, Germany Institut für Werkstoffwissenschaft, Technische Universität Dresden, Helmholtzstraße 7, D-01069 Dresden, Germany Institut für Strukturphysik, Technische Universität Dresden, Haeckelstraße 3, D-01062 Dresden, Germany
A R T I C L E I N F O
A B S T R A C T
Keywords: PMN-PT Ferroelectric Thin films Microstructure Dislocation structure Stacking faults
Pb(Mg1/3Nb2/3)O3-PbTiO3 is known for its excellent ferroelectric and electrocaloric properties. Epitaxial films, which allow applying high voltages, can be grown in a much broader parameter range when depositing them on miscut substrates than on those without miscut. To understand the influence of substrate miscut, here the structure and microstructure of two films, with and without miscut, is compared in detail. In both cases the epitaxial strain is relaxed, but while in films without miscut perfect misfit dislocations are observed, the film with miscut exhibits many partial dislocations and stacking faults. The stacking faults may play an important role for the stabilization of the pure perovskite Pb(Mg1/3Nb2/3)O3-PbTiO3 phase in thin films on miscut substrates.
1. Introduction
LaAlO3). Additionally, partial dislocations with Burgers vectors b = a / 2〈110〉 resulting in a stacking fault (SF) in-between are observed in some of these materials [12,13,14,15]. Such O-sublattice preserving stacking faults have also been reported for bulk 0.65PMN-0.35PT [16]. Whereas a significant number of microstructural data is available for simple perovskites, such detailed knowledge is lacking for PMN-PT thin films. Typically, integral X-ray diffraction measurements are used to provide evidence for the growth of a single phase material and to demonstrate the desired epitaxial relationship, whereas transmission electron microscopy (TEM) cross sections give an overview over the general microstructure. However, to our best knowledge no detailed microstructural investigations focusing on lattice defects in PMN-PT thin films were published so far. Just in reference [17] misfit dislocations are reported for 0.67PMN-0.33PT thin films grown on La0.5Sr0.5CoO3/CeO2/YSZ buffered silicon without giving further details. Therefore, our aim is to perform such detailed microstructural investigations on epitaxial 0.68PMN-0.32PT thin films grown on La0.7Sr0.3CoO3 (LSCO) buffered SrTiO3 (STO). In general, the parameter window for the growth of pure perovskite PMN-PT thin films without the lead-deficient pyrochlore phase is quite narrow, if pulsed laser deposition (PLD) is applied [21,22]. However, it was demonstrated that the growth on STO substrates with a defined
The perovskite (1 - x)Pb(Mg1/3Nb2/3)O3-xPbTiO3 (PMN-PT) shows appealing electrocaloric properties [1,2] in addition to the well-known large dielectric constants and high piezoelectric coefficients [3,4] and therefore may become a suitable material for cooling applications [5]. So far, many studies were conducted on bulk PMN-PT [6,7] as well as on thin films [8,9,10] to determine the caloric properties. Among them, thin films of PMN-PT are of great interest because of the high applicable electric fields [6] and the possibility to modify the crystal orientation and strain state through epitaxy. At the same time, a detailed microstructural analysis is required to understand the influence of the growth parameters and the resulting structure on the functional properties of the films. Therefore, we focus in this report on the defect structures of epitaxial PMN-PT films grown on SrTiO3 substrates with different miscut. Microstructural investigations on lattice defects of epitaxial ABO3 perovskites thin films typically show that these materials contain perfect misfit dislocations with Burgers vectors b = a〈100〉 to accommodate the elastic strain (for example for SrZrO3 [11], BaZrO3 [11], SrTi0.5Zr0.5O3 [11], BaTiO3 [12] and PbSc0.5Ta0.5O3 [13] grown on SrTiO3 as well as for Ba0.75Sr0.25TiO3 [14] and SrTiO3 [15] grown on
⁎
Corresponding author at: Institute for Metallic Materials, IFW Dresden, P.O. Box 270116, D-01171 Dresden, Germany. E-mail addresses: [email protected] (P. Chekhonin), [email protected] (M. Mietschke), [email protected] (D. Pohl), [email protected] (F. Schmidt), [email protected] (S. Fähler), [email protected] (W. Skrotzki), [email protected] (K. Nielsch), [email protected] (R. Hühne). http://dx.doi.org/10.1016/j.matchar.2017.05.003 Received 21 December 2016; Received in revised form 28 March 2017; Accepted 6 May 2017 Available online 07 May 2017 1044-5803/ © 2017 Elsevier Inc. All rights reserved.
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Fig. 1. XRD of PMN-PT thin films grown on LSCO buffered STO substrates: a) Θ-2Θ-scans (A10 was mounted with an inclination to compensate the 10° substrate tilt) and b) to e) pole figures of both samples. b) and d) show the (110) pole figures of STO of samples A0 and A10, respectively, while c) and e) depict (220) pole figures of PMN-PT of A0 and A10, respectively.
2.2. Instrumentation
miscut is helpful to stabilize the desired perovskite PMN-PT phase [18,19]. In this work, two pure perovskite PMN-PT thin film samples are compared, which were prepared on a STO substrate without as well as with a 10° miscut using slightly different deposition parameters. The central aim of this study is to depict microstructural differences between these two samples in order to clarify the influence of the miscut on the film growth. In the following, the perovskite PMN-PT phase is referred to just as PMN-PT. Additionally, it should be noted that the crystallographic structure of bulk 0.68PMN-0.32PT is not cubic at room temperature [20]. This composition is close to the so-called morphotropic phase boundary, where monoclinic, rhombohedral or tetragonal lattice structures coexist. However, the quantitative differences between these structures are small, so that PMN-PT will be treated as a cubic lattice with respect to crystallographic directions and planes.
The thin film samples were analyzed by X-ray diffraction (XRD). Θ2Θ-scans were performed with a Bruker D8 Advance diffractometer and pole figures were recorded in a Philips X'Pert four circle diffractometer using Co-Kα and Cu-Kα radiation, respectively. The determination of inplane lattice parameters was done by recording reciprocal space maps (RSMs) in an additional Philips X'Pert diffractometer equipped with a thin film optics using a Goebel mirror in the primary beam and narrow Soller slits on the detector side. Surface characterization was carried out with a Bruker Dimension Icon atomic force microscope (AFM). Transmission electron microscopy investigations were done in a FEI Tecnai T20 TEM equipped with a LaB6 cathode operated at 200 kV acceleration voltage and an aberration corrected FEI Titan3 80–300 TEM operated in scanning TEM (STEM) mode at 300 kV using a high angle annular dark-field (HAADF) detector. Cross-sectional TEM lamellae were prepared out of the central areas of both samples applying the focussed ion beam (FIB) technique in a FEI Helios 600i FIB and subsequent argon ion milling. The image plane in all TEM lamellae is (010). The Burgers vectors for dislocations are typically determined by a Burgers circuit in the HAADF images.
2. Experimental 2.1. Sample Preparation 0.68PMN-0.32PT thin films were deposited on single crystalline STO substrates applying pulsed laser deposition using a 30 nm to 50 nm thin LSCO buffer layer. One sample (referred to as A0) was produced on a STO substrate with the normal direction parallel to [001]. For A0 the epitaxial relationship along the film plane is: STO (001) || LSCO (001) || PMN-PT (001) and along an in-plane direction: STO [100] || LSCO [100] || PMN-PT [100]. Another sample (referred to as A10) was grown on a (001)-oriented STO substrate with a 10° miscut towards the [100] direction showing the same epitaxial relationship. The lattice parameters of bulk STO, LSCO and PMN-PT are 3.905 Å, 3.839 Å and about 4.02 Å, respectively. Therefore, in the case of epitaxy without strain relaxation, a biaxial in-plane tensile strain and biaxial in-plane compressive strain is introduced to the LSCO and PMN-PT layers, respectively. The deposition of both samples was performed in an in-house assembled PLD chamber equipped with a KrF excimer laser (248 nm wavelength, Lambda Physics) at a temperature of 550 °C. Sintered La0.7Sr0.3CoO3 and single crystalline 0.68PMN-0.32PT disks were used as targets. The deposition of sample A10 was performed at constant oxygen pressure of 0.1 mbar and a laser energy density on the target surface of about 1.5 J/cm2, more details are published in reference [19]. In contrast, sample A0 was prepared with an energy density of 0.8 J/cm2 and an oxygen pressure of 0.3 mbar to meet the narrow deposition window for standard substrates. Further details can be found in [22]. After deposition, both samples were transferred to a furnace and annealed for 1 h at 400 °C in pure oxygen at atmospheric pressure to reduce oxygen vacancies.
3. Results 3.1. X-ray Diffraction XRD Θ-2Θ-scans reveal the successful growth of LSCO and the perovskite PMN-PT phase (Fig. 1 a)) in both samples. Essentially, the (00ℓ) peaks of PMN-PT, STO and LSCO are clearly visible, whereas no sign for a pyrochlore structure was found. Some traces of non-epitaxial PMN-PT are observable as marked in Fig. 1 a). This is in accordance with the presence of a small number of protruding grains at the PMN-PT film surface (cf. one example in Fig. 4 a)) with different orientation. The AFM maps in Fig. 2 show that such grains exist in both samples, but apparently have a much higher area density in A10 in comparison to A0. However, the volume density is in general quite low. (110) pole figures of STO and (220) pole figures of PMN-PT prove the above mentioned epitaxial relationship for sample A0 (Fig. 1 b) and c)) and A10 (Fig. 1 d) and e)), respectively. Measured (110) pole figures of the LSCO layers (not shown here) are qualitatively similar to the (110) STO pole figures of the corresponding sample. RSMs of (103) reflections provide more details on the structural properties of the layers (Fig. 3). In each RSM, three distinct intensity maxima are visible, which arise from the STO, LSCO and PMN-PT layer. The in-plane lattice parameter (a-axis) of the layers, which correspond to the Qx values of the reciprocal lattice, were calculated using the (103) and (103) maxima, while the out of plane lattice parameter (caxis) was extracted from corresponding (002) peaks of the Θ-2Θ-scans. The results are listed in Table 1. 235
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Fig. 2. AFM mappings of a) A0 and b) A10. The bright spots are protruding grains.
Within the measurement uncertainty of RSM, the in-plane lattice parameter of LSCO is similar in A0 and A10 and equal to the value of STO (a = 3.905 Å), thus indicating zero or negligible strain relaxation. The c-axis of the LSCO layers shrinks as expected in the presence of biaxial in-plane tensile strain. On the other hand, the larger in-plane lattice parameter of PMN-PT indicates nearly complete strain relaxation. The slightly larger PMN-PT c-axis for both samples may be explained either by some remaining compressive in-plane strain that may be affected by oxygen vacancies or by a non-cubic crystal structure [20]. Obviously, the in-plane strain in PMN-PT is not affected by the substrate miscut. It was concluded from the RSM measurements that the [001] directions of the substrate and each subsequent layer are aligned exactly parallel in sample A0. This is not the case in A10, where the LSCO layer is slightly tilted towards the substrate normal by about 0.1° (positive tilt), while the PMN-PT layer is tilted by approximately −0.8° in the opposite direction (i.e. away from the substrate normal). These tilts might be explained qualitatively as a consequence of a stepped surface by the Nagai model [23]. In this case, the tilt angle α can be calculated by the Nagai Eq. (1)
⎛ c − cB ⎞ tan α = −⎜ E ⎟ tan φ ⎝ cB ⎠
Table 1 Lattice constants (a and c) and film thickness (t) determined by RSM, XRD and TEM, respectively, of the LSCO and PMN-PT layers. A0 aLSCO in Å (by RSM) cLSCO in Å (by Θ-2Θ) aPMN-PT in Å (by RSM) cPMN-PT in Å (by Θ-2Θ) tLSCO in nm (by TEM) tPMN-PT in nm (by TEM)
3.91 ± 3.78 ± 4.02 ± 4.03 ± ≈55 ≈160
A10 0.01 0.01 0.02 0.01
3.91 ± 3.83 ± 4.02 ± 4.06 ± ≈ 30 ≈ 300
0.01 0.01 0.02 0.01
with the measured lattice constants for sample A10 from Table 1. It is noticeable that the calculated tilt of the PMN-PT layer is lower than the measured one, which might be explained by the observed microstructure described later. It should be mentioned in this context that even an extended Nagai model, which is based on a step configuration with different step heights [24], fails to explain the higher tilt measured for sample A10.
3.2. Microstructure (1)
Cross sectional TEM bright field images with the electron beam parallel to the [010] zone axis of PMN-PT are shown in Fig. 4. The images reveal clear interfaces between the distinct layers for samples A10 (Fig. 4 a)) and A0 (Fig. 4 e) ), respectively. The layer thickness of LSCO and PMN-PT shows only minor variations in the nm range. The LSCO layer for A0 is with about 55 nm considerably thicker than
where φ is the miscut angle and cE and cB are the out-of-plane lattice parameters of the tilted epitaxial layer and the substrate or buffer layer beneath, respectively. The theoretical tilt angle towards the STO substrate normal is 0.2° for LSCO and − 0.4° (=−0.6° + 0.2° (considering the tilt of the LSCO layer beneath)) for PMN-PT using Eq. (1)
Fig. 3. RSMs of a) sample A0 and b) A10 showing the (103) reflexions of PMN-PT, LSCO and STO.
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Fig. 4. a) to c) show bright field TEM images of sample A10 at different magnifications highlighting the lamellar structure. In d) a HAADF STEM image demonstrates a stacking fault (SF), its projection is marked by a red line. e) and f) show bright field TEM images of sample A0. The [010] zone axis is always parallel to the electron beam. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
10 of these with b = a / 2[101] (Fig. 5 b)), 4 with b = a / 2[101] (Fig. 5 c)) and 4 with b = a[100]. Fig. 5 d) shows the latter case in the vicinity of a SF, which starts already in the LSCO layer. Despite the small number of identified cases, one may presume that the partial MDs at the interface are obviously unequally distributed, with b = a / 2[101] preferred to b = a / 2[101]. This is comprehensible, as the sample symmetry for a (001) plane is not granted due to the miscut. As a consequence of partial MDs at the LSCO/PMN-PT interface, SFs are introduced to the LSCO layer of sample A10. Representative HAADF images of the STO/LSCO interface shown in Fig. 6 indicate an ideal coherent growth for both samples. While the brighter La/Sr columns are clearly distinguished from the darker Co columns in the complete LSCO layer of A0, this is only the case directly at the STO/LSCO interface for sample A10. Lattice distortions are observed in the LSCO layer just a few nm above that interface. As a result, the contrast of the columns changes continuously along a chosen atomic plane. As stated above, SFs are apparently present inside the LSCO layer of sample A10, however, (partial) dislocations could not be revealed. Therefore, the SFs in the LSCO layer of sample A10 are presumably continuous and bordered by partial MDs in PMN-PT at the LSCO/PMN-PT interface. Whereas extended PMN-PT SFs start occasionally already in the LSCO layer (Fig. 5 d)), most of them appear above the LSCO/PMN-PT interface of sample A10 bordered by a partial dislocation at this position (Fig. 7 a) and b)). Five cases were identified: In three cases the Burgers vector of the partial dislocations is b = a / 2[101] and in the two other cases b = a / 2[101] and b = a / 2[101]. Some of the SFs do not extend up to the sample surface, but are bordered by another partial dislocation, two examples are shown in Fig. 7 c) and d). Here also five cases were identified, four of them with b = a / 2[101] and one with b = a / 2[101]. Finally, perfect dislocations with b = a[001] are also present in the PMN-PT layer of A10 (Fig. 7 c)). In contrast to A10, no partial MDs were observed at the LSCO/PMNPT interface for sample A0. Instead, partial dislocations very rarely interrupt the otherwise widely extended zig-zag shaped SFs in the PMNPT layer of this sample. The lattice appears to be noteworthy bend in the vicinity of SFs in
intended resulting in some cracks. However, the RSM results indicate that these cracks do not lead to a significant strain relaxation. Lamellar aligned structures are revealed in the PMN-PT layer of sample A10 (Fig. 4 b)), whereupon HAADF STEM images (Fig. 4 d)) indicate that they originate from stacking faults. Since the HAADF intensity strongly depends on the atomic number, Pb columns appear as bright spots, while columns with Mg, Nb and Ti atoms appear less bright due to the lower scattering rate resulting from their lower density. Therefore, SFs can easily be recognized by a lattice displacement along the a / 2〈101〉 direction. This is especially the case in few locations, where the normal direction of the SF is perpendicular to [010], as exemplary marked by a red line in Fig. 4 d). However, the normal direction of the SF is typically not aligned perpendicular to [010]; therefore, they appear broadened in the TEM images. The majority of the SFs in A10 have a preferred alignment. The typical angle between the SF normal direction (more accurately its projection in the (010) plane) and the thin film normal is about 40° and therefore about 30° with respect to [001], considering the 10° substrate miscut in this sample. Their alignment is never straight but seems to change direction. Also the width of the projection is different at different locations. Thus, a preferred crystallographic alignment seems to be absent. The distance between two SFs is of the order of 15 nm along the thin film normal, variations of a few nm are typical. As can be seen in Fig. 4 c), the SFs do not start immediately at the LSCO/PMN-PT interface. SFs are also present in sample A0, but here they appear in a zig-zag-arrangement (Fig. 4 f)) at the centre or the upper part of the PMN-PT layer. Throughout the thickness of the PMN-PT layer, never more than one SF is present at a certain position. Therefore, the density of SFs in A0 is about one order of magnitude lower compared to sample A10. A closer look at the LSCO/PMN-PT interface of A0 reveals that strain relaxation is realized only by misfit dislocations (MDs) with Burgers vectors b = a[100]. A typical example is shown in Fig. 5 a). In contrast, such perfect in-plane MDs are not the rule at the LSCO/PMN-PT interface of A10. Instead, the interface is dominated by partial MDs, where Fig. 5 b) is presenting a typical case. Along the LSCO/PMN-PT interface of A10 18 MDs were identified, 237
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Fig. 5. HAADF STEM images of interface dislocations at the LSCO/PMN-PT boundary in a) A0 and b) to d) A10.
(for example the Pb columns below the red line, marked by yellow arrows in Fig. 8), so that exactly at the SF the displacement is a / 2[110]. It should be noted that the STEM micrographs illustrated in this work are excerpts of large scale, high resolution HAADF images having larger regions with blurred appearance due to strain. Therefore, an exact dislocation density cannot be provided with confident statistical
many locations of both samples (Fig. 8). In the case of a SF with a lattice displacement of a / 2[101], Pb columns are expected to shift by a / 2[100] and a / 2[001]. However, if a larger lattice area is observed along a straight line, the lattice translation is not always a / 2[001] perpendicular to that line, but has a smaller value or is even zero as demonstrated by the red line crossing a SF in Fig. 8. This is not obvious, when following the lattice from one column to the next across the SF
Fig. 6. HAADF STEM images of the STO/LSCO interface of a) A10 an b) A0.
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Fig. 7. HAADF STEM images of partial dislocations in the PMN-PT layer of A10 bordering SFs.
4. Discussion If a film exceeds a certain thickness, strain relaxation may occur by MDs [25]. This is apparently the case for both PMN-PT layers investigated. In ref. [16] it was reported that (partial) dislocations in PMN-PT remain unaffected by the electron beam induced heating. Therefore it is concluded that the formation of dislocations and SFs must have taken place during film growth or during annealing at elevated temperatures. As the dislocation energy scales with b2, dissociation of perfect dislocations with Burgers vector b = a〈100〉 into two partial dislocations with b = a / 2〈110〉 is energetically equivalent. However, additional energy is stored in the SF bordered by the partial dislocations. The frequent occurrence of SFs in A10 and their irregular alignment qualitatively indicates that their specific energy seems to be low and not strongly anisotropic. The predominant presence of partial dislocations at the LSCO/PMNPT interface of A10 might be caused by a different mechanism of MD formation as for sample A0, thus resulting in the observed microstructural differences. One argument for such a difference includes the miscut in A10 that reduces the sample symmetry. Another argument considers the smaller lattice constant of LSCO and the atomic steps in A10 at its LSCO/PMN-PT interface, which lead to an additional compressive strain along [001] in the PMN-PT layer. This strain may trigger a relaxation mechanism, which preferably creates partial MDs with b = a / 2[101]. The latter argument does consider that the mechanism responsible for the formation of partial MDs in A10 cannot only be the dissociation of perfect dislocations with b = a[100] or
Fig. 8. HAADF STEM images of a SF in A10. The red dotted line demonstrates the lattice bending close to the SF. Close to the SF yellow arrows mark Pb columns that indicate a lattice displacement of a / 2[110]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
reliability without significant additional experimental effort, which was beyond the scope of our studies. However, SFs with a lattice displacement vector of a / 2〈100〉 or Burgers circuits with b = a / 2〈100〉 were not observed.
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strained thin films [31].
b = a[101], as this would lead to a balanced distribution between partial MDs with b = a / 2[101] and b = a / 2[101] at the LSCO/PMNPT interface. To observed differences between samples A0 and A10 concerning the misfit dislocation structure are not clarified in this work and require an advanced analysis focusing on the dislocation nucleation mechanism. From the experimental data gathered for A10 it can be concluded that partial dislocations bordering the extended SFs in PMN-PT do not contribute to a unilateral strain relaxation or any lattice rotation, since the Burgers vector sum of the 10 partial dislocations observed is equal to zero. However, this is not the case for the MDs at the LSCO/PMN-PT interface, where the excess of b = a / 2[101] MDs should result in a positive lattice rotation about [010]. Even though the experimental data do not provide sufficient statistical reliability, this rotation is probably overcompensated by the existence of perfect dislocations with b = a[001], so that the overall negative rotation observed by XRD can be understood. The 10 partial dislocations enclosing the SFs in A10 indicate that strain accommodation is not responsible for their existence. In general, SFs with a displacement vector of a / 2〈110〉 give rise to nonstoichiometry, if the normal direction of a SF with a displacement vector of a / 2〈110〉 in ABO3 perovskite is not aligned with 〈110〉 [26]. In this case, a SF can accommodate BO2 excess or AO deficit in ABO3. Such SFs have been observed in non-stoichiometric TiO2-rich homoepitaxial SrTiO3 thin films [27]. This leads to a theory, why the undesired lead deficient pyrochlore phase is hampered on miscut substrates: The formation of SFs is preferred due to the miscut. If the deposition parameters result in some lead deficiency, this deficiency may be accommodated by a dense network of extended SFs, thus preventing the formation of the lead deficient pyrochlore phase. Indeed, the deposition parameters for sample A0 had to be carefully adjusted to avoid the pyrochlore phase. A sample deposited on a substrate without miscut exhibits the undesired pyrochlore phase, if produced with the same deposition parameters like sample A10 [19]. This argument is supported by the fact that the observed SFs in A10 are not aligned with 〈110〉 (hence non-stoichiometric) and that the observed partial dislocations enclosing the SFs in A10 are collectively strain neutral. The preferred alignment of the SFs in A10 indicates that their formation is obviously related to the sample symmetry. Some of the SFs may be initiated by partial misfit dislocations that exist in this sample, but most of the partial dislocations bordering a SF are clearly no misfit dislocations (e.g. Fig. 7 b)). Therefore, a different mechanism must be responsible for their formation, which cannot be provided here yet. An advanced microstructural analysis of PMN-PT films with very small miscut angles or lead rich targets might be helpful to clarify this issue. It should be noted in this context that an integral X-ray photoelectron spectroscopy measurement of sample A10 published in ref. [19] did not indicate a larger lead deficit. However, smaller local deviations cannot be excluded. This work does not include measurements regarding ferroelectric and electrocaloric properties. It is worth mentioning that SFs observed in ordered bulk Pb(Sc1/2Ta1/2)O3 seem not to interact with ferroelectric domains [28]. This could also be the case here, however it may be assumed that in the presented PMN-PT thin film A10 the high density of SFs may affect the mobility of ferroelectric domain boundaries. Advanced research is required to analyze this issue in PMN-PT thin films. In both samples it was impossible to determine the precise lattice symmetry, which presumably is not cubic [20]. The mosaicity of the PMN-PT lattice made TEM diffraction experiments difficult and the SFs (especially in A10) lead to broad peak splitting in the TEM diffraction plane. Useful (large angle) convergent beam electron diffraction patterns, which might be used for lattice symmetry determination in PMN-PT (following the work of Wang et al. [29,30] in bulk PMN-PT), were impossible to record. This is mainly due to bending of the TEM lamellae caused by strain relaxation, which is known to occur in
5. Conclusions Two 0.68PMN-0.32PT thin films were produced on La0.7Sr0.3CoO3 (LSCO) buffered (001) oriented SrTiO3 substrates using pulsed laser deposition, one of them with a 10° miscut towards the [100] direction. It is shown that minor changes in substrate orientation can lead to significant differences in the microstructure. In the sample without substrate miscut, misfit dislocations have a Burgers vector b = a[100], while SFs and partial dislocations are absent in the vicinity of the LSCO/PMN-PT interface. In the centre or the upper part of its PMN-PT layer SFs are sporadically present and only a few partial dislocations were observable. In strong contrast, in the sample grown on a miscut substrate the majority of misfit dislocations are partial dislocations with b = a / 2[101] or b = a / 2[101], while there is a strain induced surplus of the former case. The microstructure of its PMN-PT layer is dominated by closely aligned SFs that are bordered by partial dislocations. Presumably, the miscut results in a favoured formation of extended non-stoichiometric SFs, whose formation is not yet fully understood. Furthermore, the authors suggest that these SFs may play an important role in the accommodation of a lead deficiency, which might originate from the preparation of PMN-PT films using non-optimized deposition parameters. Consequently, the formation of the undesired lead deficient pyrochlore phase is impeded and the perovskite PMN-PT phase is stabilized in thin films on miscut substrates. Further research is suggested to clarify this issue. Acknowledgements 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 technical support as well as A. Pöhl for preparation of the TEM lamellae. References [1] J.F. Scott, Electrocaloric materials, Annu. Rev. Mater. Res. 41 (2011) 229–240. [2] X. Moya, S. Kar-Narayan, N.D. Mathur, Caloric materials near ferroic phase transitions, Nat. Mater. 13 (2014) 439–450. [3] S.-E. Park, T.R. Shrout, Ultrahigh strain and piezoelectric behavior in relaxor based ferroelectric single crystals, J. Appl. Phys. 82 (1997) 1804–1811. [4] J. Wang, K.H. Wong, H.L.W. Chan, C.L. Choy, Composition control and electrical properties of PMN-PT thin films around the morphotropic boundary, Appl. Phys. A Mater. Sci. Process. 79 (2004) 551–556. [5] S. Fähler, U.K. Rößler, O. Kastner, J. Eckert, G. Eggeler, H. Emmerich, P. Entel, S. Müller, E. Quandt, K. Albe, Caloric effects in ferroic materials: new concepts for cooling, Adv. Eng. Mater. 14 (2012) 10–19. [6] J. Hagberg, A. Uusimäki, H. Jantunen, Electrocaloric characteristics in reactive sintered 0.87 Pb(Mg1 ∕ 3Nb2 ∕ 3)O3–0.13 PbTiO3, Appl. Phys. Lett. 92 (2008) 132909. [7] B. Rožič, M. Kosec, H. Uršič, J. Holc, B. Malič, Q.M. Zhang, R. Blinc, R. Pirc, Z. Kutnjak, Influence of the critical point on the electrocaloric response of relaxor ferroelectrics, J. Appl. Phys. 110 (2011) 064118. [8] A.S. Mischenko, Q. Zhang, R.W. Whatmore, J.F. Scott, N.D. Mathur, Giant electrocaloric effect in the thin film relaxor ferroelectric 0.9 PbMg1/3Nb2/3O3–0.1 PbTiO3 near room temperature, Appl. Phys. Lett. 89 (2006) 242912. [9] T.M. Correia, J.S. Young, R.W. Whatmore, J.F. Scott, N.D. Mathur, Q. Zhang, Investigation of the electrocaloric effect in a PbMg2/3Nb1/3O3-PbTiO3 relaxor thin film, Appl. Phys. Lett. 95 (2009) 182904. [10] D. Saranya, A.R. Chaudhuri, J. Parui, S.B. Krupanidhi, Electrocaloric effect of PMN–PT thin films near morphotropic phase boundary, Bull. Mater. Sci. 32 (2009) 259–262. [11] P.A. Langjahr, F.F. Lange, T. Wagner, M. Rühle, Lattice mismatch accommodation in perovskite films on perovskite substrates, Acta Mater. 46 (1998) 773–785. [12] T. Suzuki, Y. Nishi, M. Fujimoto, Analysis of misfit relaxation in heteroepitaxial BaTiO3 thin films, Phil. Mag. A 79 (1999) 2461–2483. [13] B.I. Birajdar, A. Chopra, M. Alexe, D. Hesse, Crystal defects and cation ordering domains in epitaxial PbSc0.5Ta0.5O3 relaxor ferroelectric thin films investigated by high-resolution transmission electron microscopy, Acta Mater. 59 (2011) 4030–4042. [14] Y.Q. Wang, W.S. Liang, P.K. Petrov, N.M. Alford, Dissociation of misfit and threading dislocations in Ba0.75Sr0.25TiO3 epitaxial film, Mater. Charact. 62 (2011) 294–297. [15] Y.L. Qin, C.L. Jia, K. Urban, J.H. Hao, X.X. Xi, Dislocations in SrTiO3 thin films
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