Residual Strain-Mediated Multiferroic Properties of Ba0.85Ca0.15Zr0

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Residual strain-mediated multiferroic properties of Ba0.85Ca0.15Zr0.9Ti0.1O3/La0.67Ca0.33MnO3 epitaxial heterostructures Songbin Li, Chuanbin Wang, Qiang Shen, and Lianmeng Zhang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b05747 • Publication Date (Web): 30 Jul 2019 Downloaded from pubs.acs.org on July 31, 2019

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ACS Applied Materials & Interfaces

Residual Strain-Mediated Multiferroic Properties of Ba0.85Ca0.15Zr0.9Ti0.1O3/La0.67Ca0.33MnO3 Epitaxial Heterostructures Songbin Li, Chuanbin Wang*, Qiang Shen, Lianmeng Zhang State Key Lab of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, China

ABSTRACT In this study, artificial multiferroic Ba0.85Ca0.15Zr0.1Ti0.9O3/La0.67Ca0.33MnO3 (BCZT/LCMO) epitaxial heterostructures were deposited on Nb-doped SrTiO3 substrates using pulsed laser deposition. The epitaxial growth of the heterostructures on the substrate was demonstrated by XRD RSM and TEM analyses, which displayed decreasing residual strain with increasing BCZT layer thickness. The electrical, magnetic and magnetoelectric properties of the epitaxial heterostructures were investigated in detail, and they were sensitive to the varying BCZT layer thickness in terms of residual strain. The multiferroic nature of the heterostructures was demonstrated by ferroelectric and ferromagnetic hysteresis loops. Strain-mediated magnetoelectric behavior was observed in the epitaxial heterostructures. Finally, dielectric properties were enhanced in the heterostructures over the BCZT single layer film, and the maximum magnetoelectric coefficient of the heterostructures was 206.5 mV/cm·Oe.

*

Corresponding author. E-mail address: [email protected]

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KEYWOEDS: BCZT/LCMO, residual strain, electrical properties, ferromagnetic properties, magnetoelectric properties 1. INTRODUCTION Multiferroic magnetoelectric (ME) materials, which are able to realize the coexistence of and the coupling effect between ferroelectricity and ferromagnetism, have been the subject of ever-increasing interest due to their potential multifunctional applications such as high sensitivity transducers, actuators and multiple-state memory1– 3

Unfortunately, although intrinsic ME behavior has been discovered in a small group

of single-phase ME materials, it is too weak at room temperature or appears at very low temperatures, which consequently limits its further study and practical applications.4 The frustration of searching for single-phase multiferroics with robust room temperature ME behavior has oriented the research to highlight artificial multiphase multiferroic

compounds

containing

ferroelectric

and

ferromagnetic

phases.

Investigations of multiphase multiferroic compounds have found a way to realize strong extrinsic ME behavior at room temperature, which is closely related to cross-coupling between the ferroelectric-ferromagnetic interfaces and ferroic ordering. In recent years, extensive work,5–7 has demonstrated that the ME coupling effect in bulk multiphase multiferroic compounds is strain mediated. Generally, the external magnetic field causes strain fluctuation in the ferromagnetic layer via the magnetostriction effect, which transports to the ferroelectric layer across the interface and in turn induces voltage variation due to the piezoelectric effect. By contrast, more than one responsible mechanism has been proposed for the ME coupling effect in


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artificial multiferroic heterostructures, such as strain-mediated,8 exchange bias coupling9 and charge modulation mechanisms.10 The investigation of artificial multiferroic heterostructure has lately been accelerated by the rapid expansion of thin film growth techniques and improved theoretical calculations.11,12 Furthermore, the design and precise control of the epitaxial structure and substrate-imposed strain would facilitate the understanding of ME physical mechanisms on the atomic scale and optimize the ME effect of the artificial multiferroic heterostructure. Therefore, the artificial multiferroic heterostructure materials show unique superiority compared to bulk artificial multiferroic composites and are promising candidates in multiferroic functional devices. In addition, it is well known that the clamping effect stemming from the substrate suppresses the magnetostriction of the ferromagnetic layer in the artificial multiferroic heterostructure. Once the substrate induced clamping effect is relaxed to a certain level, the strain mediated ME coupling effect occurs in the artificial multiferroic heterostructure. However, there are only a few works focused on the effect of clamping on the multiferroic properties of artificial multiferroic heterostructure. Here,

a

neotype

artificial

multiferroic

heterostructure

system

is

Ba0.85Ca0.15Zr0.9Ti0.1O3/La0.67Ca0.33MnO3 (BCZT/LCMO). BCZT is selected as the ferroelectric phase due to its environmentally friendly nature and excellent piezoelectric properties comparable to the toxic PZT-based materials.13 Lanthanum manganite LCMO is chosen as the ferromagnetic component of the artificial multiferroic heterostructure because of its large magnetostriction and lattice-matched parameters to perovskite ferroelectric materials.14,15 Additionally, the metallic conductivity, double



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exchange mediated ferromagnetism and colossal magnetoresistance of LCMO provides additional freedom for multifunctional devices, and it is now attracting ever-increasing attention as a ferromagnetic phase for artificial multiferroic heterostructure. Nevertheless, as far as is known, no systematic investigation of the residual strain effect on the multiferroic properties of BCZT/LCMO epitaxial heterostructures has been carried out. In this article we alter the residual strain in the epitaxial heterostructure by changing the thickness of the ferroelectric BCZT layer and keeping the ferromagnetic layer thickness unchanged. The residual strain effect on the electrical, magnetic and magnetoelectric

properties

of

artificial

multiferroic

BCZT/LCMO

epitaxial

heterostructures is systematically investigated, and they show considerable promise for multifunctional device applications. 2. EXPERIMENTAL Artificial multiferroic BCZT/LCMO bilayer heterostructures were epitaxially grown on (100)-oriented Nb-doped SrTiO3 (Nb:STO) single crystal substrates by pulsed laser deposition (KrF laser, wavelength ~ 248 nm). The bottom LCMO layers were kept at approximately 100 nm, while the top BCZT layers were 50, 100, 150 and 200 nm, defined as BL50, BL100, BL150 and BL200, respectively. During deposition of the BCZT/LCMO layers, the substrate temperature, laser frequency and substratetarget distance were kept at 650/650°C, 10/10 Hz and 40/40 mm, respectively, while the laser energy and oxygen pressure were 300/350 mJ and 10/20 Pa, respectively. A PANalytical Empyrean four-circle diffraction system was used to analyze the crystalline structure and lattice parameters of the artificial multiferroic heterostructures.


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Transmission electron microscopy (TEM, FEI Titan Themis 200) was used to characterize the cross-sectional structure. The local piezoelectric hysteresis loops were measured by piezo response force microscopy (PFM, Digital Instruments, Nanoscope IV). The ferroelectric properties were studied using a ferroelectric testing apparatus (Radiant, Premier II) at 1 kHz, while the dielectric properties were measured by an impedance analyzer at room temperature (Agilent, E4980A). The ferromagnetic hysteresis loops were texted by a superconducting quantum interference device magnetometer at 10 K (SQUID, Quantum Design). For the ME measurement, the artificial multiferroic heterostructures were placed in a Helmholtz coil that generated an ac excitation magnetic field (Hac ~ 5.7 Oe) at 1 kHz. The ME coefficient as a function of dc magnetic field (perpendicular to the polarization direction) was measured by an ME measurement system at room-temperature (SuperME, Quantum Design).

3. RESULTS AND DISCUSSION 3.1 Structure Analysis. X-ray diffraction (XRD) patterns of the artificial multiferroic bilayer heterostructures with different BCZT layer thickness are shown in Fig. 1. It can be seen that the patterns of all samples exhibit single crystal structures with no impurity phases and are highly oriented along the Nb:STO (l00) substrate reflections. The rocking curves of the BCZT layers are displayed in the inset. The full width at half-maximum (FWHM) values of the different thickness BCZT layers are collected in Table 1, which are 0.607°, 0.418°, 0.261° and 0.250° for BL50, BL100, BL150 and BL200,



respectively. Decreased FWHM values suggest that the crystallinity and out-of plane texture of the BCZT layers improve with increasing BCZT thickness. To get more structure information and determine the residual stain in artificial multiferroic BCZT/LCMO heterostructures, reciprocal space mapping (RSM) around the (002) and (103) reflections and φ scan of the (220) reflections of the BL100 heterostructure were recorded. Figs. 2 (a) and (c) display the RSM (002) and the φ scan of the (220) reflections of the BL100 heterostructure, respectively, other samples are not shown here due to similar results. The equal Qx coordinate of the scattered X-ray intensity around RSM (002) as well as the good agreement of the φ scan peak positions between the heterostructure and the Nb:STO substrate suggest that the BCZT and LCMO layers are epitaxially grown on Nb:STO substrate. Fig. 2 (b) is the RSM pattern around the asymmetric reflection of Nb:STO (103) for the BL50, BL100, BL150 and BL200 heterostructures, respectively. It can be seen that the (103) peak of the LCMO layer for all samples is aligned with the Nb:STO substrate peak, while the peak position of the BCZT layer shows varying degrees to the left of Nb:STO, and the intensity of the BCZT (103) spot increases with increasing BCZT layer thickness, revealing a strained LCMO layer and a relaxed behavior on BCZT layer for all samples. The lattice parameters of the heterostructures with different BCZT layer thickness, determined from the horizontal and vertical RSM (103) peak positions, are shown in Fig. 2 (d). It is clear that the a-axis lattice constant of the BCZT layer increases while the c-axis lattice constant decreases with increasing BCZT layer thickness, resulting in the reduction of the c/a ratio. Furthermore, the in-plane (εa) and out-of-plane (εc) misfit


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strain of the BCZT layers are given by εa=a-a0/a0 and εc=c-c0/c0, respectively, where a and c are the calculated lattice constants of the BCZT layers and a0 and c0 are the BCZT bulk values (a0=4.007 Å, c0=4.025 Å).13 The εa and εc of the heterostructures with different BCZT layer thickness are presented in Fig. 2 (e). As seen, the in-plane compression stain in BCZT layers is expressed with a negative sign and the out-ofplane tensile strain is indicated by a positive sign, and the compression of the a-axis seems to result in an elongation along the c-axis due to the lattice mismatch between the BCZT and LCMO layers. In addition, the curves in Fig. 2 (e) reveal a tendency toward a reduction in residual strain with increasing BCZT layer thickness for both the a-axis and c-axis directions. The artificial multiferroic BL100 epitaxial heterostructure was selected for TEM analysis. Fig. 3 (a) shows the crystal structure model of the heterostructure deposited on the Nb:STO substrate, explaining the LCMO as the bottom layer and BCZT as the top layer. The cross sectional images of the BCZT-LCMO and LCMO-Nb:STO interfaces recorded by scanning transmission electron microscopy (STEM) are displayed in Figs. 3 (b) and (c), respectively, and they reveal flat and sharp interfaces between the BCZT layer, LCMO layer and the Nb:STO substrate. Figs. 3 (d) and (e) show the high-resolution TEM (HRTEM) lattice images of the BCZT-LCMO and LCMO-Nb:STO interfaces, respectively. Two straight and sharp interfaces can be observed, and the BCZT layer displays heteroepitaxy aligned well with the LCMO layer while the LCMO layer displays heteroepitaxy aligned well with the Nb:STO substrate. The selected area electron diffraction (SAED) patterns recorded from the BCZT-



LCMO and LCMO-Nb:STO interfaces are presented in Figs. 3 (f) and (g), respectively. The orderly diffraction spots demonstrate good crystalline quality for the heterostructure. In addition, the diffraction spots of the BCZT and LCMO layers are almost coincident while those of LCMO layer and Nb:STO substrate are perfectly overlapping in the same order, which denotes the high quality of the epitaxial BCZT and LCMO layer (l00) preferred orientation. According to the analyses of XRD RSM and TEM, artificial multiferroic BCZT/LCMO heterostructures are epitaxially grown on Nb:STO substrate as follows:[100]BCZT//[100]LCMO//[100]Nb:STO, and the residual strain of the epitaxial heterostructures decreases with increasing BCZT layer thickness. 3.2 Electrical Properties. The typical amplitude-voltage butterfly loops and the corresponding phase-voltage hysteresis loops of BCZT/LCMO epitaxial heterostructures with different BCZT layer thickness are plotted in Figs. 4 (a)-(d), respectively. The complete ferroelectric polarization switching of all samples is demonstrated by the well-shaped butterfly loops and the 0-180° reversibly switchable piezo response phase. It is evident that the amplitude-voltage butterfly loops of the heterostructures show varying degrees of asymmetry behavior, which is also reported in previous works.16,17 This behavior can be ascribed to the built-in electric field caused by the electrode self-poling effect, which promotes downward polarization switching and in turn hinders upward polarization rotation.18 Furthermore, the effective piezoelectric coefficient d33 of the epitaxial heterostructures is determined from formula19: d33=(D-D1)/(V-V1), where D is the measured value of piezoelectric displacement, V is the corresponding applied voltage


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at each point of the amplitude-voltage loops, D1 is the piezoelectric displacement and V1 is the applied voltage of the intersection. The d33 values of all samples calculated from the amplitude-voltage loops are recorded in Table 1. Results show that the piezoelectric response of the heterostructures increases with the increase in BCZT layer thickness, which is mainly attributed to the reduction of the residual strain in the BCZT layer. Previous works20,21 have proposed that high residual stress in ferroelectric materials strengthens ferroelectric domain pinning. The pinning effect hinders polarization switching and prevents the domain wall motion, and in turn degrades the piezo response to the field-induced response. Fig. 5 (a) shows the polarization hysteresis loops of artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures. The determined corresponding remnant polarization (Pr) and electric coercive field (EC) are plotted in Fig. 5 (b). The ferroelectric nature of the artificial multiferroic heterostructures is confirmed by the typical ferroelectric hysteresis loops, whereas the symmetry and saturation of the hysteresis loops are improved with increasing BCZT layer thickness. Furthermore, an increase in Pr and a reduction in EC can also be observed when the BCZT layer thickness of the epitaxial heterostructures increases. Since the LCMO layer is a ferromagnetic phase, the polarization of the artificial multiferroic BCZT/LCMO epitaxial heterostructures should stem from the ferroelectric BCZT phase. This phenomenon may be better understood by the explanation of the passive interfacial layer effect and the reduction of leakage current with increasing BCZT layer thickness indicated in Fig. 5 (c). A voltage shift to negative bias can be observed for all



heterostructures, indicating the presence of a passive interfacial layer which is defined as a nonferroelectric layer with non-switching behavior at the BCZT-LCMO interface. Some published works on PZT/LSMO and PZT/CFO multilayer heterostructures reported a similar voltage-shift phenomenon.22,23 The source of the passive interfacial layer may be caused by the presence of misfit dislocations24 between the BCZT and LCMO layer. This could be avoided by selecting materials with high lattice matching and optimizing preparation process in future work. As pointed out in previous works,24,25 the passive interfacial layer generates screening charges in the ferroelectric layer, and strongly hinders the back switching of the domains in the meantime. Therefore, a reduction in the portion of the passive interfacial layer due to the increase in BCZT layer thickness leads to an increase in Pr and decrease in EC of the epitaxial heterostructures. Furthermore, the larger leakage current of the epitaxial heterostructures with thinner BCZT layer also leads to an inadequate poling degree with lower Pr and wider polarization hysteresis loops with larger EC. Fig. 6 (a) presents the room-temperature dielectric constant (εr) of artificial multiferroic BCZT/LCMO epitaxial heterostructures with different BCZT layer thickness, and BCZT, LCMO single layer films, respectively. All samples display similar dielectric behavior in that the εr decreases monotonically with increasing frequency. The εr values of the BL50, BL100, BL150 and BL200 epitaxial heterostructures at 1 kHz are placed in Table 1, and they are 1852, 1370, 690 and 521, respectively. Interestingly, larger εr is obtained in the artificial multiferroic BCZT/LCMO epitaxial heterostructures than the BCZT single layer films (εr ~ 387 at


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1 kHz), and the εr of the epitaxial heterostructures decreases with increasing BCZT layer thickness. The dielectric behavior found in our work is completely opposite to the related published works.26,27 This should be contributed by the high εr of the LCMO layer (εr ~ 5146 at 1 kHz). The dielectric response detected in the LCMO single layer film can be attributed to the peculiar phase-separated electronic state theory.28,29 According to this phase-separated state, LCMO is considered to be composed of a mixture of insulating and metallic ferromagnetic regimes, and the large difference in conductivity between them induces a dielectric response stemming from the spacecharge or interfacial polarization. Fig. 6 (b) displays the dielectric loss (tan δ) as a function of frequency for the epitaxial heterostructures with different BCZT layer thickness, and corresponding BCZT and LCMO single layer films. It can be observed that the tan δ of the heterostructures declines with increasing BCZT layer thickness. The larger tan δ in the epitaxial heterostructures with thinner BCZT layer thickness should be ascribed to the combined effect of (i) the high tan δ from the LCMO layer, (ii) larger leakage current and (iii) the interfacial polarization, which is more likely to activate at low frequency. 3.3 Magnetic and Magnetoelectric Properties. In the artificial multiferroic BCZT/LCMO epitaxial heterostructures with different BCZT layer thickness, typical ferromagnetic hysteresis loops were observed, as displayed in Fig. 7, which indicates the ferromagnetic nature of the epitaxial heterostructures. With the increase in BCZT layer thickness, the saturated magnetization (MS) of the epitaxial heterostructures slightly increases, while the



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coercive magnetic field remains almost unchanged. This phenomenon is similar to the results

found

in

BiFeO3/BaTiO330

and

PbZr0.52Ti0.48O3/CoFe2O431

bilayer

heterostructures, which is likely to be attributed to the ferroelectric-ferromagnetic interface effect and the strain state in the ferromagnetic layer. Regarding the calculation of the magnetization, only the volume of the LCMO layer was considered, as the fact that the ferroelectric BCZT layer has no contribution to the heterostructure in terms of magnetism. Furthermore, the strain state of the LCMO layer is almost constant with increasing BCZT layer thickness, which can be deduced by the XRD RSM analysis. Therefore, the increase in MS of the heterostructure with increasing BCZT layer thickness may be attributed to the interface effect between BCZT and LCMO layers. However, further study is necessary to clarify the effect of the BCZT layer thickness on the oxidation states of Mn in LCMO layer. Fig. 8 presents the transverse ME coefficient (αE31) dependence of the dc magnetic field (Hdc) for artificial multiferroic BCZT/LCMO epitaxial heterostructures with different BCZT layer thickness, measured by a lock-in amplifier under ac magnetic field (Hac~5.7 Oe) at 1 kHz. The value of αE31 for all samples increases monotonically with increasing Hdc, and then gets saturated until the Hdc reaches 5000-6000 Oe. The ME behavior of artificial multiferroic BCZT/LCMO epitaxial heterostructures is quite different from their bulk laminated composites32, which increases rapidly to achieve a maximum value at approximately 2000 Oe, and then decreases with further increasing magnetic field. This is caused by the clamping effect in the LCMO layer that derived from the Nb:STO substrate induced strain, which would hinder the domain wall motion.


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Consequently, the saturated magnetostriction of the LCMO layer cannot be achieved until the Hdc is as large as 5000-6000 Oe. In addition, the maximum values of αE31 for the epitaxial heterostructures obtained from the plots in Fig. 8 are assembled in Table 1, and they are 99.8, 152.3, 173.2 and 206.5 mV/cm·Oe for BL50, BL100, BL150 and BL200 epitaxial heterostructures, respectively. It is noted that the larger the residual strain, the smaller the ME coupling of the epitaxial heterostructure. Many works32-36 have confirmed that the ME coupling was mainly dominated by the magnetostrictivepiezoelectric interaction coupled via strain at the ferroelectric-ferromagnetic interfaces. Considering that the magnetostrictive effect of the LCMO layer is almost unchanged due to the same strain state (thickness), the improvement in the value of αE31 with the increasing thickness of BCZT layer could be attributed to the decrease in residual strain and the enhancement in piezoelectric effect of the BCZT layer. The observed ME coupling behavior suggests the dominant strain-mediated mechanism in the artificial multiferroic BCZT/LCMO epitaxial heterostructures. Furthermore, a significantly enhanced ME effect is achieved in the epitaxial heterostructures compared to their bulk laminated composite (αE31 ~ 6.57 mV/cm·Oe)37. In our previous work, the bulk laminated composites were sintered at high pressure together with a fast cooling rate that should cause large residual stress, and Ca elemental interdiffusion between the BCZT and LCMO phases occurred at the interface. These factors greatly degrade the ME coupling effect of the bulk laminated composite. In contrast, the BCZT and LCMO thin films were epitaxially grown on the substrates with a perfect interface structure. This achieved interface coupling at the atomic level. Consequently, the artificial



multiferroic BCZT/LCMO heterostructure obtained a larger ME effect than the bulk laminated composites. 4. CONCLUSIONS In conclusion, we deposited artificial multiferroic BCZT/LCMO heterostructures with different BCZT layer thicknesses via pulsed laser deposition. The XRD and TEM results demonstrated that the bottom LCMO layer was epitaxially grown on the substrate while the top BCZT layer was epitaxially grown on the bottom LCMO layer, and the residual strain in the epitaxial heterostructures decreased with increasing BCZT layer thickness. The well-shaped ferroelectric and ferromagnetic hysteresis loops demonstrated the multiferroic nature of the heterostructure. The multiferroic properties of the epitaxial heterostructures were sensitive to the varying BCZT layer thickness in terms of residual strain. The observed ME behavior demonstrated that a strain-mediated ME coupling mechanism dominates in the epitaxial heterostructures. In addition, enhanced dielectric properties were obtained in the heterostructure compared to the BCZT single layer, and an improved ME effect compared to the bulk laminated composites with the maximum magnetoelectric coefficient of 206.5 mV/cm·Oe has been achieved in the heterostructure, which shows promising application for new types of multiferroic functional devices.

ACKNOWLEDGEMENT This work was financially supported by the National Natural Science Foundation of China (51272195 and 51521001), the International Science and Technology Cooperation Project of Hubei Province (2016AHB008), the Natural Science


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Foundation of Hubei Province (2016CFA006) and the Fundamental Research Funds for the Central Universities (WUT: 2019Ⅲ029).

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(Ba0.7Ca0.3)TiO3/Ba(Zr0.2Ti0.8)O3

Ferroelectric

Superlattices. ACS Appl. Mater. Inter. 2015, 7, 26301–26306. (17) Luo, B. C.; Wang, D. Y.; Duan, M. M.; Li, S. Orientation-Dependent Piezoelectric Properties in Lead-Free Epitaxial 0.5BaZr0.2Ti0.8O3-0.5Ba0.7Ca0.3TiO3 Thin Films. Appl. Phys. Lett. 2013, 103, 385–991. (18) Lin, Q.; Wang, D.; Li, S.; Damjanovic, D. Strong Effect of Oxygen Partial Pressure on Electrical Properties of 0.5Ba(Zr0.2Ti0.8)O3-0.5(Ba0.7Ca0.3)TiO3 Thin Films. J. Am. Ceram. Soc. 2015, 98, 2094–2098. (19) Reddy, S. R.; Bhanu Prasad, V. V.; Bysakh, S.; Shanker, V.; Joardar, J.; Roy, S. K. Ferroelectric and Piezoelectric Properties of Ba0.85Ca0.15Ti0.90Zr0.10O3 Films in 200 nm Thickness Range. J. Am. Ceram. Soc. 2019, 102, 1277–1286. (20) Berfield, T. A.; Ong, R. J.; Payne, D. A.; Sottos, N. R. Residual Stress Effects on Piezoelectric Response of Sol-gel Derived Lead Zirconate Titanate Thin Films. J. Appl. Phys. 2007, 101, 1926. (21) Ion, V.; Craciun, F.; Scarisoreanu, N. D.; Moldovan, A.; Andrei, A.; Birjega, R.; Ghica, C.; Di Pietrantonio, F.; Cannata, D.; Benetti, M.; Dinescu, M. Impact of Thickness Variation on Structural, Dielectric and Piezoelectric Properties of (Ba,Ca)(Ti,Zr)O3 Epitaxial Thin Films. Sci. Rep. 2018, 8, 2056.



(22) Lin, R.; Wu, T. B.; Chu, Y. H. Interface Effects on the Magnetoelectric Properties of (00l)-Oriented Pb(Zr0.5Ti0.5)O3/CoFe2O4 Multilayer Thin Films. Scripta Mater. 2008, 59, 897–900. (23) Chaudhuri, A.R.; Ranjith, R.; Krupanidhi S. B.; Mangalam R. V. K.; Sundaresan, A. Interface Dominated Biferroic La0.6Sr0.4MnO3/0.7Pb(Mg1/3Nb2/3)O3-0.3PbTiO3 Epitaxial Superlattices. Appl. Phys. Lett. 2007, 90, 122902. (24) Tagantsev, A. K.; Gerra, G. Interface-Induced Phenomena in Polarization Response of Ferroelectric Thin Films. J. Appl. Phys. 2006, 100, 2623–314. (25) Jing, W.; Zheng, L.; Wang, J.; He, H.; Nan, C. Effect of Thickness on the Stress and Magnetoelectric Coupling in Bilayered Pb(Zr0.52Ti0.48)O3-CoFe2O4 Films. J. Appl. Phys. 2015, 117, 759. (26) Li, T.; Li, K.; Zhou, H. Thickness and Frequency Dependence of Magnetoelectric Effect for Epitaxial La0.7Sr0.3MnO3/BaTiO3 Bilayer. J. Alloys Compd. 2014, 592, 266– 270. (27) Dai, Y. Q.; Dai, J. M.; Tang, X. W.; Zhang, K. J.; Zhu, X. B.; Yang, J.; Sun, Y. P. Thickness Effect on the Properties of BaTiO3-CoFe2O4 Multilayer Thin Films Prepared by Chemical Solution Deposition. J. Alloys Compd. 2014, 587, 681–687. (28) Rivas, J.; Mira, J.; Rivasmurias, B.; Fondado, A.; Dec, J.; Kleemann, W.; Rodriguez, M. A. Magnetic Field-Dependent Dielectric Constant in La2/3Ca1/3MnO3. Appl. Phys. Lett. 2006, 88, 55. (29) Dagotto, E.; Hotta, T.; Moreo, A. Colossal Magnetoresistant Materials: the Key Role of Phase Separation. Phys. Rep. 2001, 344, 1–153.


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Zhao, X.; Nan, C.


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W. Wang, A. P.; Li, H. B. Strain-Mediated Converse Magnetoelectric Coupling in La0.7Sr0.3MnO3/Pb(Mg1/3Nb2/3)O3-PbTiO3

Multiferroic

Heterostructures.

Cryst.

Growth Des. 2018, 18, 5934−5939. (37) Li, S. B.; Wang, C. B.; Shen, Q.; Hu, M. Z.; Zhang, L. M. Thickness Ratio Effect on Multiferroic Properties of BCZT-LCMO Laminated Composites Prepared by Plasma Activated Sintering. J. Alloys Compd. 2018, 762, 415–421.

Table 1 FWMH and measured properties of artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures. BCZT thickness BL50 BL100 BL150 BL200

FWMH (°) 0.607 0.418 0.261 0.250

d33+ (pm/V) 40±5 51±5 61±5 91±5

d33(pm/V) 51±5 60±5 78±5 89±5


εr 1852 1370 690 521

αE31 (mV/cm·Oe) 99.8 152.3 173.2 206.5

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Fig. 1 XRD patterns of artificial multiferroic BCZT/LCMO heterostructures with different BCZT layer thickness. The inset shows the rocking curves of the corresponding BCZT layers.

Fig. 2 (a) RSM (002) of the BL100 heterostructure. (b) RSM (103) of the BL50, BL100, BL150 and BL200 heterostructures, respectively. (c) XRD φ scan of the BL100 heterostructure. (d) The a, c-axis lattice constant and c/a ratio of heterostructures with different BCZT layer thickness. (e) The in-plane (εa) and out-of-plane (εc) residual strain of heterostructures with different BCZT layer thickness.



Fig. 3 Crystal structure model of an artificial multiferroic BCZT/LCMO heterostructure grown on Nb:STO substrate (a). Cross sectional HAADF-STEM images of the BCZTLCMO (b) and LCMO-Nb:STO (c) interfaces. HRTEM images of the BCZT-LCMO (d) and LCMO-Nb:STO (e) interfaces. SAED patterns of the BCZT-LCMO (f) and LCMO-Nb:STO (g) interfaces.

Fig. 4 Local piezo response amplitude voltage butterfly loops and phase voltage


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hysteresis loops for (a) BL50, (b) BL100, (c) BL150 and (d) BL200 epitaxial heterostructures.

Fig. 5 (a) Ferroelectric hysteresis loops of the artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures. (b) The remnant polarization and coercive field of the epitaxial heterostructures with different BCZT layer thickness. (c) Voltage dependence of leakage current density for epitaxial heterostructures with different BCZT layer thickness.

Fig. 6 Frequency dependence of dielectric constant (a) and loss tangent (b) for artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures as well as corresponding BCZT and LCMO single layer films.



Fig. 7 Ferromagnetic hysteresis loops of the artificial multiferroic heterostructures with different BCZT layer thickness.

Fig. 8 Magnetic field dependence of the ME coefficients for artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures.


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Table of Contents graphic



Fig. 1 XRD patterns of artificial multiferroic BCZT/LCMO heterostructures with different BCZT layer thickness. The inset shows the rocking curves of the corresponding BCZT layers. 217x168mm (300 x 300 DPI)


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Fig. 2 (a) RSM (002) of the BL100 heterostructure. (b) RSM (103) of the BL50, BL100, BL150 and BL200 heterostructures, respectively. (c) XRD φ scan of the BL100 heterostructure. (d) The a, c-axis lattice constant and c/a ratio of heterostructures with different BCZT layer thickness. (e) The in-plane (εa) and outof-plane (εc) residual strain of heterostructures with different BCZT layer thickness. 399x241mm (300 x 300 DPI)



Fig. 3 Crystal structure model of an artificial multiferroic BCZT/LCMO heterostructure grown on Nb:STO substrate (a). Cross sectional HAADF-STEM images of the BCZT-LCMO (b) and LCMO-Nb:STO (c) interfaces. HRTEM images of the BCZT-LCMO (d) and LCMO-Nb:STO (e) interfaces. SAED patterns of the BCZT-LCMO (f) and LCMO-Nb:STO (g) interfaces. 279x147mm (300 x 300 DPI)


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Fig. 4 Local piezo response amplitude voltage butterfly loops and phase voltage hysteresis loops for (a) BL50, (b) BL100, (c) BL150 and (d) BL200 epitaxial heterostructures. 483x359mm (300 x 300 DPI)



Fig. 5 (a) Ferroelectric hysteresis loops of the artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures. (b) The remnant polarization and coercive field of the epitaxial heterostructures with different BCZT layer thickness. (c) Voltage dependence of leakage current density for epitaxial heterostructures with different BCZT layer thickness.


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Fig. 6 Frequency dependence of dielectric constant (a) and loss tangent (b) for artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures as well as corresponding BCZT and LCMO single layer films.



Fig. 7 Ferromagnetic hysteresis loops of the artificial multiferroic heterostructures with different BCZT layer thickness. 222x170mm (300 x 300 DPI)


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Fig. 8 Magnetic field dependence of the ME coefficients for artificial multiferroic BL50, BL100, BL150 and BL200 epitaxial heterostructures. 213x171mm (300 x 300 DPI)



Table of Contents graphic 471x276mm (300 x 300 DPI)


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Residual Strain-Mediated Multiferroic Properties of Ba0.85Ca0.15Zr0

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