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Ferroelectric phase transition induced a large FMR tuning in self-assembled BaTiO3:Y3Fe5O12 multiferroic composites Guohua Dong, Ziyao Zhou, Xu Xue, Yijun Zhang, Bin Peng, Mengmeng Guan, Shishun Zhao, Zhongqiang Hu, Wei Ren, Zuo-Guang Ye, and Ming Liu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b06876 • Publication Date (Web): 15 Aug 2017 Downloaded from http://pubs.acs.org on August 17, 2017

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ACS Applied Materials & Interfaces

Ferroelectric phase transition induced a large FMR tuning in self-assembled BaTiO3:Y3Fe5O12 multiferroic composites

Guohua Dong†, Ziyao Zhou*†, Xu Xue†, Yijun Zhang†, Bin Peng†, Mengmeng Guan†, Shishun Zhao†, Zhongqiang Hu†, Wei Ren†, Zuo-Guang Ye‡, Ming Liu*† †

Electronic Materials Research Laboratory, Key Laboratory of the Ministry of

Education & International Center for Dielectric Research, Xi’an Jiaotong University, Xi’an 710049, China ‡

Department of Chemistry and 4D LABS, Simon Fraser University, Burnaby, British

Columbia, V5A 1S6, Canada

KEYWORDS: magnetoelectric coupling, vertical aligned heterostructure, YIG, BTO, multiferroics

ABSTRACT: Yttrium iron garnet (YIG) is of great importance in RF/microwave devices for its low loss, low intrinsic damping and high permeability. Nevertheless, tuning of YIG based multiferroics is still a challenge due to its near zero magnetostriction and the difficulty of building epitaxial interface between ferromagnetic garnet and ferroelectric perovskite phases. In this work, the vertical aligned heterostructure of YIG:BTO/STO(001) with local epitaxial interface between BTO and YIG is well-constructed, where the single crystal BTO pillars are embedded 1

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in YIG matrix. A large magnetoelectric coupling effect that drives YIG’s FMR shift up to 512 Oe and 333 Oe (1-2 order greater than all states of the art progresses) is obtained through BTO ferroelectric phase changes induced by temperature variation at 295 K and 193 K, correspondingly. This record high magnetoelectric tunability of YIG paves a way towards thermal/electrical tunable YIG devices.

1. INTRODUCTION The central challenge in multiferroic field is to achieve large magnetoelectric (ME) tunability while maintaining small linewidth, low loss and high resistivity, particularly, in ferrite based multiferroics. A great deal of work has been demonstrated with giant ME coupling tunability upto ~600-800 Oe in ferrite based multiferroics,1-3 nevertheless, these spinel ferrites such as Ni ferrite, Co ferrite and Mn ferrite have relatively big linewidth (>300 Oe), limiting their application in GHz frequency devices. Yttrium iron garnet (Y3Fe5O12, YIG) ferrite with very low intrinsic damping (α ~ 10-5)4,5, high Curie temperature (TC ~ 550 K)6,7 and large band gap (Eg ~ 2.85 eV)8,9, which has overwhelming advantages over most of ferrites, is widely used in microwave filters, magneto-optical applications and acoustic transmitters,10-12 especially for rapidly growing communications industry behind mobile radio and satellite communications, as well as radar and other microwave technologies. Furthermore, voltage controllable YIG based multiferroics through strain/stress ME coupling without noisy, bulky and energy consuming electromagnets are of great significance for realizing next generation RF/microwave devices that are low loss, 2


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compact, fast and energy efficient.13,14 For example, an electric field-tunable YIG-PZT microwave resonator had been reported with a ferromagnetic resonance (FMR) frequency shift of 18-25 MHz for an electric field E=10 kV/cm, where the FMR tunable range is very limited.15 Other YIG based multiferroic heterostructures like gadolinium gallium garnet (GGG)-YIG-Pt-BSTO-Pt multilayers were built, where researchers had achieved 5 Oe FMR field shift by applying 25 V onto BSTO layer, resulting in a small tunability of 2 MHz at 9.5 GHz.4 The small tunability in YIG-based multiferroics may come from several reasons: firstly, most of the integrated ME devices are constructed by bonding the YIG and piezoelectric phases together with glue, which greatly degrades the strain/stress transfer efficiency at the poor interface;16 secondly, it is extremely difficult to grow epitaxial YIG directly onto ferroelectric substrate for large lattice mismatch;17 thirdly, YIG has near zero magnetostriction constant;18 fourthly, the substrate clamping effect also reduces the ME tunability in multiferroic heterostructures.19 In order to overcome this challenge, other multiferroic structures beside thin film heterostructures were proposed. The multiferroic composites can be classified into the following three groups: lamellar multilayer structure (2-2 type), vertically aligned heteroepitaxial structure (1-3 type) and nanoparticles in matrix type (0-3 type).20,21 Among these types, 1-3 multiferroic composites exhibit a much higher ME coupling coefficient for its negligible clamping constraints from the substrate and the larger interfacial area between the two phases.22 Vertical aligned nanocomposite multiferroic thin films have been established and investigated intensively since the pioneering 3



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works by H. Zheng, H. Wang and Q. X. Jia groups. Particularly, spinel ferrimagnetic (such as: CoFe2O4, NiFe2O4 and MgFe2O4) pillars embedded in a ferroelectric BaTiO3 or

BiFeO3

(BTO,

BFO)

matrix

are

typical

examples

for

two-phase

magnetoelectrically coupled systems. The abundant effective strain, phase and interface couplings between the magnetostrictive and piezoelectric phases at the vertical interfaces have been well-studied by the theoretic calculation and plenty of experimental results,23-28 leading to enormous enhancements of ME coupling effects. 1-3 type vertical aligned multiferroic structure opens a door for voltage tunable YIG multiferroics, however, few garnet and perovskite vertically aligned composite structures have been reported because high quality YIG crystals typically require a high temperature synthesis process. Meanwhile, the chemical and crystalline compatibility between YIG and ferroelectric perovskite phase is very poor. In this work, we have initialed to construct the self-assembled YIG garnet and BTO perovskite vertically aligned composite structure, where BTO is the most extensively investigated ferroelectric material that is widely used in composite multiferroic materials. We prepared the YIG:BTO heterostructure onto STO(001) substrate under 850°C by pulse laser deposition (PLD) at an oxygen pressure of 200 mTorr with a repetition rate of 3 Hz. After annealing at 1200 °C for 30 mins, single crystal, cubic shape of BTO nano-pillars were emerged from YIG matrix, confirmed by scanning transmission electron microscopy (STEM) and energy dispersive spectrometer (EDS), which behaved differently from the majority of spinel and perovskite composite films. Most interestingly, local epitaxial interfaces between 4


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ferromagnetic garnet and ferroelectric perovskite phases were discovered and studied by atomic-scale STEM. It is well known that the bulk BTO crystals exhibits three distinct

structural

phase

transitions:

cubic-tetragonal

(C-T)

at

~393

K,

tetragonal-orthorhombic (T-O) at ~278 K, and orthorhombic-rhombohedral (O-R) at ~190 K, respectively. Large ME coupling effects are expected around these phase transition points in BTO-based composites, which originate from BTO lattice changes during phase transitions and the high strain transfer efficiency across these interfaces.29,30 In this work, ferromagnetic resonance (FMR) spectroscopy is a more direct and sensitive technique to investigate the magnetic response, especially in RF/microwave devices. As expected, large ME coupling that drives FMR shift up to 512 Oe and 333 Oe (almost 10 times larger than epoxy glued YIG/PMNPT31) is obtained through BTO phase transition at 295 K and 193 K, respectively, breaking the state of the art YIG tunability with almost an order of magnitude. It is worthy to mention that there are slight phase transition temperature differences between reported bulk BTO and BTO nanostructures in this experiment, which also exists in other nano-structural BTO phases.32 Giant ME coupling coefficients during transition are calculated as 256 Oe/K and 166 Oe/K, correspondingly. Furthermore, it is worthy to admit that the YIG presents polycrystalline behavior with large linewidth, indicating a further work to obtain single crystal YIG:BTO heterostructure with perfect FMR linewidth. This novel YIG:BTO vertical local epitaxial growth multiferroic heterostructure enables the tunable YIG based microwave devices, where the FMR can be modulated by small thermal variation. In addition, this work also 5



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paves a way toward voltage tunable YIG devices and fully epitaxial growth YIG/ferroelectric nanostructures are expected.

2. RESULTS AND DISCUSSION 2.1 Phase Identification and Structural Characterizations Phase composition studied by X-ray diffraction are shown in Figure 1(a). The pattern contains three sets of diffraction peaks which can be assigned to perovskite BTO, to garnet YIG and to the STO substrate, respectively. The BTO phase of composite film has (00l)/(h00) oriented structure with high crystallinity, while YIG shows random distribution orientations in contrast. The four peaks of the YIG phases can be identified according to the powder diffraction database (PDF#77-1998). The large lattice mismatch of YIG and STO and the high growth temperature of YIG are still the impenetrable obstacles, leading to polycrystalline structure of YIG phases.33,34 Although the quality of YIG is limited, it is worthy to mention that our goal is focused on realizing the ME coupling between YIG and BTO phases rather than establishing the perfect YIG film. Additionally, there is no other impurity or intermediate phases (e.g. YFeO3 and YBa2Fe3O8) in YIG:BTO composite film, which indicates the absence of unwanted chemical reaction between the two phases during the high-temperature annealing process. It is noted that the diffraction peaks of BTO (200)/(002) are separated, which implies a tetragonal crystal symmetry and the existence of both a and c domains at room temperature. However, XRD can only provide the macroscopic phase and structure information, as shown in Figure 1(a). 6


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Morphology and microstructures of YIG:BTO will be discussed further below.

Figure 1. Phase and surface micrographs, (a) Room temperature HRXRD pattern of a YIG:BTO composite film; The inset shows the amplification of XRD near 2θ=46° (b) Surface morphology by using SEM shows BTO rectangular pillars in the YIG matrix of the composite film; (c) Magnification of the surface morphology. Figure 1(b) shows top view SEM image of YIG:BTO composite film. Interestingly, these rectangular pillars are uniformly inserted in the matrix. The projected dimension of rectangular pillars is distributed from 500 nm to 1µm. Simultaneously, the edges of pillars exhibit well-alignment along the / directions of STO substrate. The specific shape and orientation of rectangular pillars may represent BTO phase and the other grainy phase should be YIG corresponding to the XRD results. Figure 1(c) presents the enlarged image in order to obtain precisely structural relationship between YIG and BTO. 7



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Figure 2. Self-assembled vertically aligned composite structure, (a) SAED pattern showing the epitaxy growth of the BTO on STO film in the system. The inset shows the corresponding cross-sectional TEM image. (b) A typical cross-sectional STEM image showing epitaxial BTO embedded in the YIG matrix. (c-f) Various elemental EDS mapping of the sample showing element distributions. Epitaxial grown BTO phase on STO substrates can also be identified by selected-area diffraction pattern along [100], as shown in Figure 2(a). The distinct diffraction dots from BTO and YIG suggest a sharp phase separation. The inset shows the corresponding cross-sectional bright-field TEM image. Clearly, the pillar with a diameter of ~300 nm is vertically embedded in the film with the thickness of ~220 nm. Cross-sectional STEM image of the YIG:BTO film is shown in Figure 2(b). The bright contrast pillars are combined with the dark film, which provide a direct interpretation of different phases in the high angle annular dark field (HAADF) imaging mode. Meanwhile, these pillars are grown straight from substrate to the above film surface, indicating a highly oriented growth. These EDS elemental maps for Ba/Ti, Y, Fe and integrated color map from the cross-sectional STEM image are 8


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shown in Figure 2(c-f), respectively. The small energy difference between BaLα and TiKα is about 42 eV, therefore, their spectral lines always overlap each other, as shown in Figure 2c. The presence of separated BTO pillar and YIG matrix phase is clearly demonstrated by STEM and EDS mapping, meanwhile these chemical analyses verifies the well-separated YIG:BTO phases without diffusion reaction. 2.2 Microstructure Characterizations of Vertically Aligned YIG:BTO film These pillar-matrix interfaces are crucial in composite films, especially for the ferromagnetic and ferroelectric coupling system, where the ME coupling effects at these interfaces strongly depend on the interfacial conditions.35,36 Cross-section STEM-HAADF image of a selected BTO pillar within the YIG matrix is demonstrated in Figure 3(a) and its chemical composition is identified by EDS mapping, as shown in Figure S1. The sharp distribution of elements (Ba, Y and Fe) further confirm that BTO and YIG phase are absolutely separated in the composite. Figure 3(b-c) show atomic resolution HAADF images taken at the BTO/YIG/STO, BTO/STO and YIG/STO interfaces, respectively. Figure 3(b) presents a STEM image at the triple junction where both YIG and BTO phases meet the substrate. The ambiguous region between the two phases can be attributed to a structural rearrangement due to the lattice mismatch between BTO and YIG. Surprisingly, there exists a certain mathematic matching relationship between the two phases (3dYIG(332)=2aBTO(100)=8.01 Å) at the interface from these atomic resolution STEM-HAADF images, as shown in Figure 3(d), although BTO and YIG are hardly combined due to the large lattice mismatch and incompatible crystal chemistries. In 9



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other words, the crystal structure of YIG:BTO shows the local epitaxial growth between BTO and YIG at the interface. During annealing process, BTO (001) had a better wettability with STO (001) substrates, however, YIG with weak thermodynamic partially disturbed the initial nucleation of BTO and therefore limited the lateral growth of BTO. Ultimately, BTO pillars embedded in YIG matrix are emerged in the YIG:BTO composite film. The SEM image (Figure 1(b)) also provides further evidence for the existence of the ordering distribution of these BTO pillars. Moreover, the interface of YIG and STO is similar with the BTO/YIG. The clearly high resolution STEM image for the interface of YIG and STO is shown in Figure S2(b). The construction of hetero-interface in YIG:BTO films and local epitaxial multiferroic heterostructure may provide a possibility for achieving strong ME coupling. In addition, as shown in Figure 3(c), the sharp interface between BTO and STO is obtained and Fourier-filtered image of STEM image for the interface shows Moire patterns, presenting the periodicity of misfit dislocation cores observed along the interface. The atomic resolution STEM image about interface between BTO and STO is presented in Figure S2(a), which can provide a sharp interface by each atom columns. The misfit dislocations cause partial strain relaxation associated with the lattice mismatch between the BTO and STO.

10


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Figure 3. Nano-structural characteristics of YIG:BTO film, (a) Cross-sectional STEM-HAADF image of YIG:BTO /STO heterostructures in low magnification. (b) A magnified atomic-scale STEM image illustrating the triple junction for BTO, YIG and STO. (c) An atomic-scale STEM image showing the epitaxial hetero-interface between BTO and STO. The inset shows Fourier-filtered image of STEM images for the local interface (d) Schematic illustration of crystallographic model for BTO/YIG/STO triple junction interface. 2.3 Magnetic Properties of YIG:BTO Composite Films The magnetic properties of YIG:BTO composite film were investigated, by applying H-field along two directions, as shown in Figure 4(a). The hysteresis loops show that the easy axis of magnetization of the film is parallel to the film plane due to shape anisotropy.9 For the inset of Figure 4(a), soft magnetic properties with 27.5 Oe of coercivity along in-plane (IP) and 124.5 Oe of coercivity along out-of-plane (OOP) are observed, accordingly. These results are comparable with YIG full films.37,38 The 11



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polar graph in Figure 4(b) shows the angle dependence of the FMR field for YIG:BTO composite film at room temperature. The in-plane and out-of-plane resonance field are 2783 and 4071 Oe, respectively, which indicate a weak in-plane anisotropy. As shown in Figure 4(c), the magnetization loops in the temperature range from 50 to 350 K are summarized along IP by using SQUID, showing soft magnetic properties.

Figure 4. Magnetic characteristics of YIG:BTO film, (a) Magnetic hysteresis loops of YIG:BTO film measured at room temperature in both IP and OOP. The inset shows the M-H loops near the zero field. (b) The angle dependence of the FMR resonance field for YIG:BTO composite film at room temperature at 9.2 GHz. (c) Magnetic hysteresis loops of YIG:BTO film measured at 50 K, 150 K, 250 K and 350 K, respectively. The inset shows the M-H loops near the zero field. (d) Magnetization-Temperature curves recorded in FC and ZFC modes for the YIG:BTO film in an external magnetic field of 50 Oe The inset of Figure 4(c) shows a magnified region of hysteresis loops around the zero field. It is observed that the coercive field is larger at lower temperature and this is a typical temperature dependent magnetic behavior of a ferromagnetic material. 12


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Figure 4(d) shows the temperature dependence of the dc susceptibility measured under a field of 50 Oe and by performing both zero-field-cooled (ZFC) and field-cooled (FC) conditions. The FC curve runs above the ZFC one, as the magnetization originating from room temperature induction with a weak external magnetic field (50 Oe). FC and ZFC results show the presence of ferromagnetic phase, which gives a decrease of magnetization with increasing temperature (the blocking temperature is near 400 K) and the irreversibility between the ZFC and FC magnetizations.39 2.4 Ferroelectric Phase Transition Induced a Large ME Coupling Figure 5(a) shows the temperature dependent FMR field of YIG:BTO composite film measured at 9.2 GHz with the external magnetic field parallel and perpendicular to the film, respectively. The FMR can be expressed by Kittel equations for both in-plane and

out-of-plane

cases.

In-plane: = γ(H + H )(H + H + 4πM ) ;

Out-of-plane: = γ(H + H , − 4πM ), where f is frequency in megahertz, γ is the gyromagnetic ratio, Hr is the resonance field supplied by an external electromagnet, and Heff is the in-plane effective magnetic anisotropy. In this case, MS decreases with the temperature increasing, and Hr for IP will increase, while it will decrease for OOP. More interestingly, there are two mutation regions of in-plane FMR field with decreasing temperature of YIG:BTO composite film. BTO crystallizes in a perovskite structure undergoes three structural phase transitions on cooling process. As shown in the Figure 5(a), the FMR field measurements exhibit two characteristic jump-drop regions at 295K and 193 K, which corresponding to the phase transitions 13



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of BTO from the T-O and the O-R, respectively.29,40 The two phase transition mutations can be attributed to the structural phase transitions inducing magnetoelastic coupling at the interface between BTO and YIG. As for strain mediated ME coupling, the interface quality is very critical to establish the ME response between ferroelectric and ferromagnetic phases. STEM data confirm the existence of local epitaxial growth between BTO and YIG, which plays a great important role on transferring stress. Meanwhile, the strength of the in-plane ME coupling is larger than out-of-plane. The foremost reason is that the BTO pillars have about twice as height as YIG matrix from STEM and SEM images, which gives rise to larger (tensile or compressive) stress along the in-plane direction during the BTO phase transition. These BTO pillars consisting of two possible domain orientations, a-domains and c-domains, should be taken into account at the T phase. With a/c-domains reorientations during the BTO phase transition, the two kinds of stress (tensile and compressive) act on the YIG matrix, which may give rise to the dynamic peculiarity of jump-drop for the trend of the FMR field. Schematic representation of a/c-domains of the BTO unit cells near the O/T phase boundary are shown in Figure 5(b). BTO is weakly tetragonal (in the bulk lattice parameters are a=3.992 Å and c=4.035 Å at room temperature) and orthorhombic phase (O-phase) appears with decreasing temperature (lattice parameters are a=3.980, b=4.010 and c=4.020). During the BTO phase transition, the polarization will go to reorientation and the tensile strain of ~1.06% along the in-plane direction and compressive strain of ~ 0.1% are calculated along out-of-plane direction. These primary calculations assume the equal percentage of a domain and c 14


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domain, therefore, resulting in a much larger FMR field shift for in-plane than that of the out-of-plane. It is worthy to re-emphasize that this FMR field shift (depending on the temperature) in BTO:YIG is originated from the strain mediated ME coupling effect. The large strain change from temperature induced by BTO phase transition drives the magnetic anisotropy of YIG. Here, temperature is only a controlled parameter and the key mechanism is the ME coupling between BTO and YIG phases. Because it is very challengeable to apply electric field (E-field) directly onto each separated BTO pillars, then we study the strain induced magnetism change via temperature driven BTO phase transition. Therefore, this over 500 Oe FMR field shift of YIG is essentially a strain mediated ME effect, not temperature effect.

Figure 5. ME coupling effects in YIG:BTO film, (a) The FMR field dependence of temperature in both IP and OOP. (b) Schematic representation of a/c domain of the BTO unit cells near the O/T phase boundary. (c) Contour plot of FMR dependence of 15



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a narrow temperature range from 302 K to 289 K, showing FMR behavior during phase transition. Figure 5(c), as contour plot of FMR phase diagram, shows the FMR field dependence of a narrow temperature range from 302 K to 289 K with 1 or 0.5 K step size, where the color represents the FMR intensity. A giant FMR field change of 512 Oe occurs among 2 K near the phase transition (corresponding to a giant ME coupling coefficient of 256 Oe/K), while the intensity of FMR peaks is greatly enhanced. The FMR intensity is related to the FMR microwave absorption. During BTO phase transition, the dielectric constant of BTO will change significantly and lead to a corresponding microwave absorption change. This phenomenon has been also found in the Fe3O4/PZN–PT heterostructure.41 Furthermore, the IP and OOP FMR linewidths of the YIG:BTO as a function of temperature are shown in Figure S3. Although these FMR linewidth changes are relatively week along both IP and OOP directions, the FMR linewidth near the phase transition clearly decreases. Meanwhile, the linewidth of YIG:BTO gradually increases with decreasing temperature due to the inhomogeneous broadening. Additionally, the temperature dependence of FMR spectra of YIG:BTO composite film near 193 K from the in-plane is shown in Figure S4, which corresponds to the Figure 5c. Simultaneously, it is demonstrated that O-R phase transition stress of BTO can tune FMR field of YIG (A FMR field change of 333 Oe). This tunability of YIG is much larger than the state of the art reports4, 31, 42, where the substrate clamping constraint effect may lead to a very poor ME coupling strength. In contrast, for self-assembled YIG:BTO composite film, the local epitaxial lattice match between BTO pillars and YIG matrix provide a much greater ME 16


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coupling effect, which does not suffer from substrate clamping effect. Lastly, for our system, the FMR field becomes back to normal level when deviating from the phase transition temperature of BTO, which can be attributed to relaxation behavior of the crystallographically induced strain.40

3. CONCLUSIONS Self-assembled vertical aligned YIG:BTO composite films are deposited on STO (001) substrates by PLD from mixed YIG:BTO composite target. After annealing, rectangular BTO pillars are uniformly inserted in the YIG matrix. The crystal structure and chemical analysis demonstrate that perfect epitaxial interfaces are formed among BTO pillars, YIG matrix and STO substrates. Giant ME coupling induced FMR field shifts of 512 Oe and 333 Oe (1-2 order greater than that of the state of the art YIG based multiferroics) are achieved at 295 K and 193 K, corresponding to the T-O and O-R phase transitions of BTO, respectively. This novel YIG:BTO nanostructures build a foundation of realizing voltage/thermal tunable YIG based RF/microwave devices.

4. EXPERIMENTAL SECTION Sample fabrication: Self-assembled vertical aligned YIG:BTO nano-composite films were grown by the pulsed laser deposition technique in a vacuum chamber with a base pressure of approximately 2 × 10-7 Torr at a temperature of 850 °C by using KrF excimer laser (λ = 248 nm). The composite target containing 65 mol% of BTO 17



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and 35 mol% of YIG is prepared by a conventional ceramic sintering method. The stoichiometric mixture of high-purity BaCO3, TiO2, Y2O3, and Fe2O3 powders were ground, pressed, and then sintered at 1200 °C for 4 h to synthesize BTO and YIG powders respectively. The calcined BTO powders were then mixed with YIG powders and then pressed into a disk and subsequently sintered at 1250 °C for 3 h to make the composite targets. These composite films were deposited at an optimized oxygen pressure of 200 mTorr with a repetition rate of 3 Hz onto STO (001) substrates. After deposition, the composite thin films were cooled in 200 Torr oxygen at a cooling rate of 5 °C/min. In order to obtain the superior magnetic properties and good crystallinity, the as-grown composite films were subsequently annealed at an optimized temperature of 1200 °C for 30 mins in oxygen atmosphere to get well-built structure and performance. Crystal structure and phase characterization: The phase and nanostructure of the YIG:BTO composite films were characterized by high-resolution X-ray diffraction (HRXRD, PANalytical X'Per MRD) and TEM. The cross-sectional specimens were prepared through cutting, manual grinding and dimpling. The specimen thickness was brought down to a residual value of 15 to 20 µm, and final thinning of the specimen was performed on a Gatan PIPS. For high-resolution STEM and EDS mapping, a JEOL ARM200F microscope with an aberration corrector for a probe-forming system was operated at 200 kV. Magnetic properties characterization: The M-H hysteresis loops of YIG:BTO composite films on STO substrates were investigated by vibrating sample 18


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magnetometer (VSM, Lake Shore 7404) and superconducting quantum interference device (SQUID). FMR spectra measurements were performed by an X-band (≈9.2 GHz) electron paramagnetic resonance system (JEOL, JES-FA200) operated at TE 011 mode. The microwave power is maintained at 1 mW in this experiment.

Supporting Information The Supporting Information is available free of charge on the ACS Publications website. Figure S1: The chemical composition of YIG:BTO composite film. S2: The atomic resolution STEM images for the interface of BTO/STO and YIG/ STO, respectively. S3: The IP and OOP of FMR linewidth for the YIG:BTO as a function of temperature. S4: The temperature dependence of FMR spectra of YIG:BTO composite film near 193K from the in-plane.

AUTHOR INFORMATION

Corresponding Author *E-mail: [email protected]; [email protected] Notes The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The work was supported by the Natural Science Foundation of China (Grant Nos. 51472199, 11534015, and 51602244), the Natural Science Foundation of Shaanxi Province (Grant No. 2015JM5196), the National 111 Project of China (B14040), the 973 Program (Grant No. 2015CB057402), and the Fundamental Research Funds for the Central Universities. The authors appreciate support from the International Joint Laboratory for Micro/Nano Manufacturing and Measurement Technologies. Z.-G. Y. acknowledges support from the Natural Sciences and Engineering Research Council of Canada (NSERC, Grant No.203773). Z.Z. and M.L. are supported by the China Recruitment Program of Global Youth Experts.

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