Room-Temperature Multiferroics and Thermal Conductivity of 0.85

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Room Temperature Multiferroics and Thermal Conductivity of 0.85BiFe TiMgO-0.15CaTiO Epitaxial Thin Films (x = 0.1 and 0.2) 1-2x

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Ji Zhang, Wei Sun, Jiangtao Zhao, Lei Sun, Lei Li, Xue-Jun Yan, Ke Wang, ZhengBin Gu, Zhenlin Luo, Yanbin Chen, Guoliang Yuan, Ming-Hui Lu, and Shantao Zhang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b06961 • Publication Date (Web): 12 Jul 2017 Downloaded from http://pubs.acs.org on July 13, 2017

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Room Temperature Multiferroics and Thermal Conductivity Conductivity of 0.85BiFe1-2xTixMgxO3-0.15CaTiO3 Epitaxial Thin Films Films (x = 0.1 and 0.2) Ji Zhang, † Wei Sun, ‡ Jiangtao Zhao, § Lei Sun, † Lei Li,† Xue-Jun Yan,† Ke Wang, ‡ Zheng-Bin Gu,† Zhen-Lin Luo, § Yanbin Chen,|| Guo-Liang Yuan,¶ Ming-Hui Lu,† Shan-Tao Zhang,†, *) † National Laboratory of Solid State Microstructures and Department of Materials Science and Engineering, College of

Engineering and Applied Science & Collaborative Innovation Center of Advanced Microstructures, Nanjing University, Nanjing 210093, China ‡ State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua

University, Beijing 100084, China § National Synchrotron Radiation Laboratory and CAS Key Laboratory of Materials for Energy Conversion, University

of Science and Technology of China, Hefei, Anhui 230026, China ||

National Laboratory of Solid State Microstructures and Department of Physics & Collaborative Innovation Center of Advanced Microstructures, Nanjing University, Nanjing 210093, China ¶

School of Materials Science and Engineering, Nanjing University of Science and Technology, Nanjing 210094, China

ABSTRACT: Thin films of 0.85BiFe1-2xTixMgxO3-0.15CaTiO3 (x = 0.1 and 0.2, abbreviated as C-1 and C-2) have been fabricated on (001) SrTiO3 substrate with and without conductive La0.7Sr0.3MnO3 buffer layer. The x-ray θ-2θ and φ scans, atomic force microscopy and cross-section transmission electron microscopy confirm the (001) epitaxial nature of the thin films with very high growth quality. Both the C-1 and C-2 thin films show well-shaped magnetization-magnetic field hysteresis at room temperature, with enhanced switchable magnetization values of 145.3 emu/cc and 42.5 emu/cc, respectively. The polarization-electric loops and piezoresponse force microscopy measurements confirm the room temperature ferroelectric nature of both films. However, the C-1 films illustrate relatively weak ferroelectric behavior and the poled states are easy to relax whereas the C-2 films show relatively better ferroelectric behavior with stable poled states. More interestingly, the room temperature thermal conductivity of C-1 and C-2 films are measured to be 1.10 W/(m·K) and 0.77 W/(m·K) respectively. These self-consistent multiferroic property and thermal conductivity are discussed by considering the compositiondependent content and migration of Fe-induced electrons and/or charged point defects. This paper not only provides multifunctional materials with excellent room temperature magnetic, ferroelectric and thermal conductivity properties, but also may stimulate further work to develop BiFeO3-based materials with unusual multifunctional properties. KEYWORDS: thin films, TEM, magnetic properties, ferroelectric nature, thermal conductivity

INTRODUCTION Multifunctional materials have the ability to simultaneously response to various external magnetic, electric, heat, optical stimuli, etc. Thus, they have plenty of applications in modern electronic devices like sensor, actuator, and memory. For such applications, room temperature integrated multifunctional properties are important since most devices are operated near room temperature. As a typical multifunctional material, multiferroics can response to magnetic and electric stimuli simultaneously. In the past decades, multiferroics has attracted intensive attention not only due to the possible applications in new devices, but also the fascinating basic physics like electric and magnetic coupling.1-4 However, it is still a challenge to develop single phase room temperature multiferroics with long-range ferroelectric and ferromagnetic orders because in general, ferroelectric order requires closed-shell d0 or s2 cations whereas ferromagnetic order favors open shell dn with unpaired electrons.5 Perovskite BiFeO3 (BFO) is one of

the very few known single phase multiferroic materials with coexisted ferroelectric (Curie temperature Tc ~ 1103 K) and canted G-type antiferromagnetic (Néel temperature TN ~ 643 K) orders above room temperature.1 Though the canted antiferromagnetic order gives rise to weak ferromagnetism, it can only induce very weak magnetic-electric coupling even in low temperature range. Up to now, the weak magnetism and weak magnetic-electric coupling is still the main obstacle preventing actual applications of BFO.6 It is well known that by using strain effect in heterostructures, single phase thin films can show some unusual behaviors different from the corresponding bulk counterparts, e.g., weak room temperature multiferroic properties have been achieved in BFO epitaxial thin films.7,8 The responsible mechanisms for unexpected properties in heterostructures are attributed to the lattice mismatch between film/substrate, the changed composition/crystal structure, the interface-induced strain, and the couplings between charges, spin, and orbital degrees of freedom, and broken symmetry at the interfaces.9

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As to BFO thin films, it is established that epitaxially strained BFO thin films have coexisted rhombohedral and tetragonal crystal structures to form strain-induced morphotropic phase boundary (MPB) and thus have enhanced ferroelectric property, the ferroelectric polarization can be larger than 100

µC/cm2.10-17 Based on numerous works on BFO thin films, room temperature ferroelectric property and the ferroelectric domain configurations of BFO have been widely investigated.10-17 However, the room temperature magnetic property of BFO is not well unexplored, and much works are still necessary to understand and improve the room temperature magnetic property. Although strain in thin films can suppress the long-range spin cycloid in BFO and enhance magnetic property to some extent,18,19 the enhancement of magnetic property by strain is still too weak to meet the requirements for actual applications. Actually, instead of using strain effect in heterostructures, chemical doping based on atomically designing is a more feasible and effective way to enhance magnetic property of BFO. Some significant progresses have been achieved in this field very recently, such as the reported BFO-based room temperature multiferroic bulks.6,20 Especially, based on the concepts of designing magnetic percolating exchange pathway and composition-induced MPB, (1-y)BiFe1-2xTixMgxO3-yCaTiO3 bulk system shows enhanced switchable polarization and magnetization at room temperature.6 As it is described above, strain in heterostructures has positive effects on macroscopic ferroelectric property11-17 and positive, though limited, effects on magnetic property,18-21 it is worthy to fabricate and investigate the multiferroic property of this system in thin film form. On the other hand, in multiferroic materials both the domain configurations and the couplings among lattice, spin, and charge degrees of freedom have direct effects on heat transport behaviors.22-25 For example, by engineering domain structure and domain density, the thermal conductivity of ferroelectric materials can be artificially tuned.24,25 Accordingly, works on thermal conductivity of multiferroic materials will be helpful for understanding the multiferroic interactions and domains. Up to date, reports on thermal conductivity of BFO and BFO-based multiferroic materials are very rare in literatures.22,26,27 So, it is interesting to investigate the thermal conductivity of BFO-based films in order to achieve more insights into multiferroic thin films. Based on the above descriptions, c-axis epitaxial thin films of 0.85BiFe1-2xTixMgxO3-0.15CaTiO3 with x = 0.1 and 0.2 (abbreviated as C-1 and C-2 respectively hereafter) have been fabricated and investigated. Enhanced, but significantly composition-dependent, room temperature multiferroic properties and thermal conductivities have been observed and discussed in detail.

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situ annealed in chamber for 10 minutes under the oxygen pressure of 0.1mbar, cooled to 400 oC with the decreasing temperature rate of 10 oC/min, then cooled to room temperature naturally. The structures of the films were investigated by x-ray diffractometer (Rigaku Ultima III and Shanghai Synchrotron Radiation Facility (SSRF)) and transmission electron microscopy (TEM, FEI Tecnai F20). For the ferroelectric measurements, about 200 nm thick Ag layer patterned with circles of around 50 nm diameter was deposited on the films as the top electrodes and LSMO served as the bottom electrode. Then the films were annealed at 550oC with rapid increasing and decreasing temperature rate of 10 oC/min. The polarizationelectric field (P-E) loops were measured at 10 kHz using a ferroelectric test unit (Precision Premier II, Radiant Techologies, USA). The surface morphology, out-of-plane (OP) phase, in-plane (IP) phase, and amplitude image were studied by using a commercial atomic force microscopy (AFM, MFP-3D, Asylum Research) with a piezoresponse force microcopy (PFM). A Pt/Ir-coated Arrow EFM cantilevers (Nanoworld, nominal spring constant 2.8 N/m, resonant frequency 75 kHz) were used. The magnetic property was measured by using a superconductor quantum interference device (SQUID, Quantum Design, MPMS-3). The thermal conductivity was measured by a home-made time-domain thermoreflectance (TDTR) method, the thermal conductivity was calculated by fitting the ratio of input and output signals against delay time.29

RESULTS AND DISCUSSION

EXPERIMENTAL SECTIONS The thin films were prepared on (001) SrTiO3 (STO) bare substrates and La0.7Sr0.3MnO3 (LSMO) buffered (001) STO substrates by pulsed laser deposition (PLD) using a 248 nm KrF excimer laser. Stoichiometric and LSMO ceramics Bi-rich (10% excess) C-1 and C-2 ceramics were used as targets for deposition.28 Before each deposition, the chamber was evacuated to a base pressure of 10-7 mbar. During each deposition for LSMO, the substrate temperature was 800oC, the laser repetition rate was 2Hz, the laser energy density was 1.2 J/cm2, and the flowing oxygen pressure was 0.1mbar. During the depositions of C-1 and C-2 films, the deposition temperature is 700oC, while other conditions are the same with the deposition for LSMO. After each deposition, the films were in

Figure 1. The x-ray θ-2θ scans of (a) C-1/STO, (b) C1/LSMO/STO, (c) C-2/STO, and (d) C-2/LSMO/STO heterostructures. The sharp diffractions peaks indicate that the C-1, C-2, and LSMO films are well crystallized with c-axis orientation. Figure 1 illustrates the local XRD θ-2θ scans of the C-1 and C-2 films on bare and LSMO buffered STO substrates. As

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can be seen, only (002) sharp diffractions peaks from C-1 and C-2 films, LSMO buffer layers and STO substrates can be detected. The weak peaks around 2θ~42° is attributed to Cu Kβ diffraction. Therefore, single phase C-1, C-2 and LSMO films are well crystallized with c-axis texture. The c-axis epitaxial nature of the C-1, C-2 and LSMO thin films is further confirmed by x-ray φ scans. The typical φ scan results of (112) reflections of C-1, LSMO, and STO of C-1/LSMO/STO heterostructures are plotted in Figure 2(a)-2(c), respectively. Clearly, four equal spaced peaks separated by 90° can be observed, as is expected for cubic/pseudocubic crystal structure. Without LSMO buffer layer, similar results can be obtained (Figure S1 in the supplementary material), which indicates that C-1 and C-2 films can be c-axis epitaxially deposited on STO. Typical AFM surface morphologies of the C-1 and C-2 films on STO substrates without LSMO buffer layers are illustrated in Figure S2 in the supplementary material. As can be seen, both films show atomically smooth surfaces without any cracks or voids. The root mean square surface roughness is as low as 262.2 pm and 138.9 pm for C-1 and C-2, respectively. Such high quality surfaces will reduce any rough surface ambiguities during the following PFM measurements, as will be shown below.

shows the selected area electron diffraction (SAED) images of the film taken along the [001] zone-axis. The slight split peak at the high-index reflections further confirms the epitaxial nature.

Figure 3. (a) The low magnification cross-section TEM images of C-1 films. HRTEM images of the interfaces for STO and LSMO (b), LSMO and BFO-CTO (c), respectively. The inset is SAED images of the film taken along the [001] zone-axis.

Figure 2. The typical x-ray φ scans on the C-1/LSMO/STO heterostructures on (112) reflections of (a) C-1, (b) LSMO, and (c) STO substrates. These results confirm the c-axis epitaxial nature of C-1 and LSMO films. Similar results are also obtained on other heterostructures. The typical dark field TEM image of C-1/LSMO/STO heterostructure is shown in Figure 3(a). As can be seen, the film has very flat surface with the thickness of 140nm, indicating the high quality of the films, which is consistent with the results of AFM. Figure 3(b)-3(c) show the high resolution transmission electron microscopy (HRTEM) images of the interfaces of LSMO/STO and C1/LSMO, respectively. Both interfaces, indicated by the white arrows, are sharp. The coherent arrangements of atoms indicate the well epitaxial growth of the LSMO and C1 films. The inset of Figure 3(c)

Figure 4. Room temperature magnetization-magnetic field (M-H) hysteresis loops of the C-1/STO and C-2/STO heterostructures. The inset shows the partly enlarged curves. Both the thin films show well shaped ferromagnetic hysteresis with switchable magnetizations. Room temperature magnetization-magnetic field (M-H) curves were measured on C-1/STO and C-2/STO heterostructures without LSMO buffer layers, the results are plotted in Figure 4. The inset shows the partly enlarged figure. Two important features should be noted. First, both the C-1 and C-2 films show ferromagnetic-like well-saturated M-H hysteresis loops with switchable magnetization. The maximum magnetization (Ms) is 145.3 emu/cc (~1.38 µB/Fe) and 42.5 emu/cc (~0.54 µB/Fe) for C-1 and C-2 films, respectively.21 By further



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considering that the theoretical density of BFO is ~8.37g/cm3,30 it is clear that our data are significantly larger than the corresponding C-1 and C-2 ceramic bulks,6 undoped epitaxial BFO films,31 and nearly 2 times larger than heavily Mn-doped BFO epitaxial films.21 Second, the Ms of the C-1 films is ~3 times larger than that of the C-2 films. The possible reason is that on the one hand, Fe cations in C-1 occupy 68% B-sites whereas in C-2, Fe cations occupy 51%, which means C-1 composition is easier than C-2 composition to form more percolating exchange pathways spanning the sample, and thus shows larger magnetization. On the other hand, Vacancies play an important role in determining the multiferroic properties of BiFeO3 based films. Cation vacancies are likely to be formed under oxygen-rich condition while oxygen vacancies under oxygen-poor condition.32 Our films are deposited under low oxygen pressure of 0.1mbar, this means the oxygen vacancies should be the main vacancy type in our cases and their contributions to multiferroic property should be comparable in the C-1 and C-2 films. However, either the magnetic property (Fig.4) or the ferroelectric property (Fig. 5, Fig. S3) differ obviously. That means the predominant effects determining the multiferroic should be the composition, i.e., the C-1 films has higher content Fe element and thus show stronger magnetic property than the C-2 films. On the other hand, though the contribution of defects such as oxygen vacancies on magnetic property is weak, but cannot be ignored completely. Oxygen vacancy, acting as donor, can provide electrons to change the valance of Fe3+ to Fe2+. The higher content Fe in C1 means the more opportunities for valence changing between Fe3+ and Fe2+, this process is accompanied with the hopping of electrons and defects, which may enhance the magnetic interaction between neighboring Fe cations by forming doubleexchange like couplings or even contribute to macroscopic magnetic behavior directly due to vacancy-induced ferromagnetism.33,34

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represented in Figure 5(a), 5(b), and Figure 6(a), 6(b), respectively. Clearly, there are observable contrast between the polarization up and polarization down ferroelectric domains, such “box-in-box” domain architectures indicate the room temperature ferroelectric nature with switchable polarization for both thin films. Furthermore, the contrast between the polarization up and polarization down of the C-2 films is much sharper than that of the C-1 films, and the negative amplitude of the C-2 films (Figure 6(c)) is obviously larger than that of the C-1 films (Figure 5(c)), both indicating that the C-2 films may show more robust room temperature switchable ferroelectric polarization, as will be discussed below. The local ferroelectric switching PFM amplitudes and phases of the C-1 and C-2 films are shown in Figure 5(c), 5(d), 6(c) and 6(d). Again, the butterfly shaped amplitude curves and nearly 180°changed hysteresis phase confirm the room temperature ferroelectric nature with switchable ferroelectric polarization, especially for the C-2 films. Actually, such room temperature ferroelectric nature of both films is further confirmed by the P-E hysteresis loops measured at 10 kHz, as shown in the Figure S3 in the supplementary material. Such high measuring frequency is helpful to decrease, though it is difficult to avoid, the effects of charged defects on P-E loop.

Figure 6. The PFM out-of-plane amplitudes phase images (ab) and the corresponding local PFM switching spectra of amplitude and phase (c-d) of the C-2 films.

Figure 5. The PFM out-of-plane amplitudes phase images (ab) and the corresponding local PFM switching spectra of amplitude and phase (c-d) of the C-1 films. Room temperature ferroelectric nature of the C-1 and C-2 films has been confirmed by PFM measurements on C1/LSMO/STO and C-2/LSMO/STO heterostructures. Both films show homogeneous ferroelectric property. Actually, the PFM measurements on different local areas are almost the same, with the error less than 8%. The typical out-of-plane (OP) PFM amplitudes and phase images of ferroelectric domains written on the C-1 and C-2 films by a ±20 V tip bias are

By comparing Figure 5, Figure 6, and Figure S3, it is found that the C-2 films manifest more stable piezoresponse than the C-1 films. The possible reasons are attributed to the following two aspects. On the one hand, compared with the C-1 composition, the C-2 composition is much closer to the substitution-induced MPB region,6 thus may have more stabilized ferroelectric domains. On the other hand, it is believed that electrons and charged defects like oxygen vacancy can migrate under electric field to form a built-in depolarization field, such a field will lead to quick disappearance or relaxation of the poled states in PFM measurements.35 In our cases, the C-2 films have relatively lower content Fe cation than the C-1 films, which means the content of electrons and charged defects resulting from valence change between Fe3+ and Fe2+, should be lower, thus the detrimental effect of electrons and charged oxygen vacancy on domain switching in PFM measurement is relatively lower, in the C-2 films. The room temperature thermal conductivity κ was measured by the TDTR method on the C-1/LSMO/STO and C2/LSMO/STO heterostructures.27,29 The measuring schematic


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is shown in the inset of Figure 7. According to the ratio of input and output signals against delay time shown in Figure 7, the room temperature thermal conductivity κ is calculated to be 1.10 W/(m⋅k) and 0.77 W/(m⋅k) for the C-1 and C-2 films, respectively. It is noticed that both the thermal conductivity values are comparable with that of multiferroic BFO thin films and other ferroelectric thin films or single crystals like Pb(Zr,Ti)O3 and Pb(Mg,Nb)O3-PbTiO3,24-27 but a little lower than that of multiferroic BFO and BFO-based ceramic bulks.22 In general, multiferroic BFO bulk usually shows larger leaky than thin film counterpart, because bulk may have more electrons and charged defects resulted from the valence change between Fe3+ and Fe2+, such electrons and charged defects can migrate in the materials, such migrating process will enhance heat flowing, thus lead to relatively larger thermal conductivity. Actually, this is the case in our experiments: The C-1 films show larger thermal conductivity than C-2 films, which can be mainly attributed to the relatively higher content of electrons and charged defects in the C-1 films, as discussed above. In one word, the migrating of such electrons and charged defects will not only enhance magnetic property and weaken ferroelectricity by accelerating the relaxation of poled states as discussed above, but also can assist transferring heat, i.e., improve thermal conductivity behavior. It should be noted that the ferroelectric domain configurations can also affect the thermal conductivity because the domain walls can act as scatter centers to reduce thermal conductivity, i.e., different domain configurations lead to different thermal conductivity.26,27 Though we cannot exclude the effects of domain on the observed different thermal conductivity, in our opinion, the effects of domain configuration on thermal conductivity cannot be responsible for the significantly different thermal conductivity in our cases since domain configuration-induced difference of thermal conductivity is very limited.27 At last, we emphasize that the thermal conductivity is homogeneous, measurements carried out at different local areas indicate an error limit of 14%.

changing can be detected, which is reasonable since the remanent magnetization of the single phase BFO-based films is very weaker, and thus the ME effect in single phase BFO-based films is very weaker, than that of composite films.36, 37

CONCLUSIONS In conclusion, c-axis epitaxial 0.85BiFe1-2xTixMgxO30.15CaTiO3 (x = 0.1 and 0.2) thin films were fabricated on bare and LSMO buffered (001) STO substrates. Both films show enhanced but significantly composition-dependent room temperature multiferroic properties with switchable magnetization and polarization. The composition dependent magnetizations and stability of ferroelectric poled states are discussed by considering the migrating of composition-induced charged defects and electrons, which is further confirmed by the room temperature thermal conductivity measurements. Our work may provide further insight into and stimulate further works on multifunctional BFO-based materials.

ASSOCIATED CONTENT Supporting Information. This material is available free of charge via the Internet at http://pubs.acs.org. X-ray φ scan results on the C-1/STO heterostructures, indicating that C-1 films are c-axis epitaxially grown on bare STO substrates; AFM surface morphologies of the (a) C-1/STO, (b) C-2/STO heterostructures, the root mean square surface roughness is 262.2 pm and138.9 pm, respectively; The room temperature polarizationelectric (P-E) loops of C-1 and C-2 films.

AUTHOR INFORMATION Corresponding Author *[email protected]

Notes The authors declare no competing financial interests.

ACKNOWLEDGMENT This work was supported by the 973 programe (2015CB921203, 2013CB632900), the National Nature Science Foundation of China (U1432112, 11374010), National key research and development program of China (2016YFA0300102), and “Dengfeng B” project of Nanjing University. Z. Luo acknowledges the beamtime and support provided by beamline 14B of SSRF.

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Figure 7. The delay time dependent heat transport curves for the C-1 and C-2 thin films, the corresponding thermal conductivity is calculated to be 1.10 and 0.77 W/(m·K), respectively. The inset shows the measuring schematic. Compared with the C-2 films, the C-1 films have better room temperature magnetic property and comparable ferroelectric property, therefore, the C-1 films were chosen as the object to probe the possible magnetoelectric (ME) coupling effect by measuring the d33 with an applied magnetic field.36 However, no magnetism-induced piezoelectric coefficient (d33)

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