Enhanced Self-Biased Magnetoelectric Coupling in Laser-Annealed

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Enhanced self-biased magnetoelectric coupling in laser annealed Pb(Zr,Ti)O3 thick film deposited on Ni foil Haribabu Palneedi, Deepam Maurya, Liwei D. Geng, Hyun-Cheol Song, GeonTae Hwang, Mahesh Peddigari, Venkateswarlu Annapureddy, Kyung Song, Yoon Seok Oh, Su-Chul Yang, Yu U. Wang, Shashank Priya, and Jungho Ryu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b16706 • Publication Date (Web): 08 Jan 2018 Downloaded from http://pubs.acs.org on January 8, 2018

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ACS Applied Materials & Interfaces

Enhanced self-biased magnetoelectric coupling in laser annealed Pb(Zr,Ti)O3 thick film deposited on Ni foil Haribabu Palneedi,† Deepam Maurya,‡ Liwei D. Geng,§ Hyun-Cheol Song,‡,

¶

Geon-Tae

Hwang,† Mahesh Peddigari,† Venkateswarlu Annapureddy,‡‡ Kyung Song,∇ Yoon Seok Oh,ǁ SuChul Yang,‼ Yu U. Wang,§ Shashank Priya,‡,* Jungho Ryu†,*

†Functional

Ceramics Group, Korea Institute of Materials Science (KIMS), Changwon 51508,

Korea ‡

Bio-inspired Materials and Devices Laboratory (BMDL), Center for Energy Harvesting Materials and Systems (CEHMS), Virginia Tech, Blacksburg, Virginia 24061, USA §

Department of Materials Science and Engineering, Michigan Technological University, Houghton, MI 49931, USA

¶ Center

for Electronic Materials, Korea Institute of Science and Technology (KIST), Seoul 02792, Korea ‡‡CSIR-National

Physical Laboratory, Dr. K.S. Krishnan Road, New Delhi 110012, India

∇Department

of Materials Modeling and Characterization, Korea Institute of Materials Science (KIMS), Changwon 51508, Korea

ǁ

Department of Physics, Ulsan National Institute of Science and Technology (UNIST), Ulsan 44919, Korea ‼

Department of Chemical Engineering, Dong-A University, Busan 49315, Korea

1 Environment ACS Paragon Plus

ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ABSTRACT Enhanced and self-biased magnetoelectric (ME) coupling is demonstrated in a laminate heterostructure comprising of 4 µm-thick Pb(Zr,Ti)O3 (PZT) film deposited on 50 µm-thick flexible Nickel (Ni) foil. A unique fabrication approach, combining room temperature deposition of PZT film by granule spray in vacuum (GSV) process and localized thermal treatment of the film by laser radiation, is utilized. This approach addresses the challenges in integrating ceramic films on metal substrates, which is often limited by the interfacial chemical reactions occurring at high processing temperatures. Laser induced crystallinity improvement in the PZT thick film led to enhanced dielectric, ferroelectric and magnetoelectric properties of the PZT/Ni composite. A high self-biased ME response on the order of 3.15 V/cm·Oe was obtained from the laser annealed PZT/Ni film heterostructure. This value corresponds to a ~2000% increment from the ME response (0.16 V/cm·Oe) measured from the as-deposited PZT/Ni sample. This result is also one of the highest reported value among similar ME composite systems. The tunability of selfbiased ME coupling in PZT/Ni composite has been found to be related to the demagnetization field in Ni, strain mismatch between PZT and Ni, and flexural moment of the laminate structure. Phase-field model provides quantitative insight into these factors and illustrates their contributions toward the observed self-biased ME response. The results present a viable pathway towards designing and integrating ME components for new generation of miniaturized tunable electronic devices. Key words: Pb(Zr,Ti)O3; Nickel; laser annealing; self-biased; magnetoelectric


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1. INTRODUCTION Magnetoelectric (ME) composites comprising of magnetostrictive and piezoelectric materials, are promising candidates for a variety of applications including magnetic sensors, voltage tunable inductors, data storage elements, spintronics, energy harvesters, etc.1-5 The ME coupling in a laminate composite, quantified in terms of ME voltage coefficient (αME), depends primarily on the interfacial strain transfer between the magnetostrictive and piezoelectric layers.6 Magnetostrictive strain (λij) exhibits quadratic variation with the external magnetic field (H) applied. In the absence of DC bias (Hbias), the magnetically induced mechanical deformation (or AC magnetostriction) is negligible in most of the magnetostrictive materials. Consequently, the piezomagnetic coefficient (qij = dλij/dH) could be nearly zero for Hbias = 0. The variation of αME is similar to that of qij, and thus an optimum DC bias is required, in addition to AC magnetic field, to achieve the maximum in ME coupling of the composites.7,

8

This necessitates the use of

permanent magnets or other DC magnetic source around ME composites, resulting in bulky devices with additional electromagnetic interference, which hinders the on-chip integration. To circumvent above mentioned issues, researches have focused on the development of self-biased magnetoelectric (SME) composites exhibiting finite αME at the zero magnetic bias.9,10 In composites, the SME coupling can be realized by inducing a built-in magnetic bias in the magnetostrictive layer. Among different approaches proposed for achieving SME, the most elegant technique relies on taking advantage of magnetization hysteresis in magnetostrictive material.10 The remanent magnetization in hysteretic magnetostrictive material can be modulated by applying AC magnetic field (shifting of the hysteresis loop on field axis) to produce an internal magnetic bias, and thereby finite remanent αME at zero bias.11 SME response has been modeled by taking into account the nature of magnetization, its tunability through varying



demagnetization state, and the resultant differential magnetic flux density distribution in the magnetostrictive layer.12, 13 For a finite non-spherical ferromagnet, the demagnetization factor has been observed to be dependent on its structural parameters (size and geometry). Optimization of these parameters results in effective control over the position and magnitude of maximum αME of the composite.14-16 However, a comprehensive understanding on the key factors influencing the SME behavior in composites is still lacking. Nickel has been the most utilized material in SME composites, based on hysteretic magnetostrictive effect.17-20 However, in majority of the studies, the ME composites were prepared by epoxy bonding of the piezoelectric layer with Ni layer, which has limited relevance for device fabrication. Few studies have pursued the deposition of piezoelectric films on Ni, employing techniques that were complex, limited to small area, and required careful optimization of deposition parameters and material compositions.21, 22 Besides, it is also necessary to optimize the post-deposition thermal treatment, adopted to enhance the properties of the piezoelectric films, for minimizing the thermal expansion mismatch between the film and substrate and to mitigate the interfacial reactions during fabrication. This study provides a fundamental understanding of the SME phenomenon in a piezoelectric Pb(Zr,Ti)O3 (PZT) thick film deposited on a flexible magnetostrictive Ni foil. The choice of PZT was based on its high piezoelectric voltage constant (gij), because, αME ∝ gij.23 Generally, the crystallization of PZT films requires thermal annealing at temperatures above 600°C.24,

25

However, such high temperatures could lead to oxidation of Ni and interfacial

chemical reactions between PZT and Ni. Although use of buffer layers between PZT and Ni and post-deposition treatment by rapid thermal annealing (RTA) could improve the mechanical adhesion and crystallization and inhibit interfacial reactions, sometimes, these approaches have


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detrimental effect on the interfacial strain transfer and ME coupling.21, 22, 26, 27 In this work, direct deposition of PZT film (4 µm-thick) was carried out using granule spray in vacuum process (GSV) at room temperature (RT).28, 29 Crystallization of the PZT film was induced through laser annealing (Figure 1), which addressed the challenges in achieving interfacial compatibility of the film with metal substrates. By doing so, we obtained a highly enhanced SME response of 3.15 V/cm·Oe in PZT/Ni composite, which is the highest magnitude among similar ME composite systems reported in literature. Experimental results on SME response in the PZT/Ni composite are explained by modeling the mesoscale phenomenon and its tunability.

2. EXPERIMENTAL METHODS

Fabrication of the ME composite: A 4 µm-thick PZT film was deposited on a thin flexible Ni foil (50 µm-thick) using Granule Spray in Vacuum (GSV) process at room temperature, followed by localized annealing of the film with continuous-wave 560 nm Ytterbium fiber laser radiation. PZT granules (d50 of primary particle and granule were ~1.3 µm and ~ 100 µm, respectively, JA-1, JK Precision Electric, Korea) mixed with medical grade-dried air were sprayed (230 L/min flow rate) on to the Ni substrate (99.5 % metals basis, Alfa Aesar, USA) through a laval type nozzle (400 mm-slit length) in a deposition chamber under vacuum (~ 4 Torr). The details of the GSV process and the stages involved in film growth are described in our prior studies.30, 31 Figure 1 illustrates the scheme of laser annealing of the PZT film. During annealing, the laser beam with an effective diameter of 50 µm was focused and illuminated over the target area (10 × 5 mm2) of the sample



mounted on an X-Y linear stage. The incident laser power and the sample scanning speed were fixed at 965 mW and 0.03 mm/s, respectively.

Figure 1. (a) and (b) Schematic illustration of the laser irradiation of the PZT film on Ni substrate and the laser annealing sequence, respectively. Laser spot size, power, scan speed, and annealed area were 50 µm, 965 mW, 0.03 mm/s, and 10 × 5 mm2, respectively. Both the deposition and laser treatment of the PZT film were conducted at room temperature.

Characterization of the ME composite: The as-deposited (AD) and laser annealed (LA) PZT films on Ni were characterized by Xray diffraction (XRD, D/Max 2200, Rigaku Corporation, Japan), Raman spectroscopy (LabRam HR800, Horiba Ltd., Japan), and transmission electron microscopy (TEM, JEM-2100F, JEOL Ltd., Japan), to understand the structural differences between them. The cross-sectional features of the PZT/Ni and the elemental distributions near the interfacial regions were examined by TEM and Energy-dispersive X-ray spectroscopy (EDS) equipped with it. The magnetization of the AD and LA PZT/Ni samples was evaluated along the in-plane direction using a physical


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properties measurement system (PPMS, Quantum Design, USA). For electrical and ME characterization, patterned circular Pt top electrodes (0.5 mm diameter) were sputter deposited on the PZT film while the Ni substrate had served as the bottom electrode. Dielectric properties of the films were measured by an impedance analyzer (4294A, Agilent Technologies, USA) and the polarization hysteresis behavior was evaluated by a ferroelectric test system (Precision LC II, Radiant Technologies, USA). For ME measurement, the samples were corona poled at 130 °C for 20 min using a DC potential of 12 kV. The ME output voltage from the PZT/Ni composite was measured by lock-in amplifier (SR-850, Stanford Research Systems, USA) under offresonance condition. The ME signal from the sample was obtained in its thickness direction (transverse αME) by subjecting it to superimposed AC (Hac = 1 Oe, f =1 kHz) and DC (Hdc) magnetic fields along its in-plane direction. Finite element modeling: Magnetic flux density distribution of Ni was estimated by a finite element model using COMSOL Multiphysics 5.2. In this model, rectangular Ni sheets with different thickness (t = 50, 100, 200 µm) but same planar dimension of 10×5 mm2 (identical to that used in experiments) were considered. A magnetostatic insulating boundary condition was applied around the Ni sheet, which is assumed to have a relative permeability of 600 when placed in air and subjected to a magnetic field (Hac) of 1 Oe along its in-plane direction.12, 13 In each case, the Ni sheet was meshed with 20000 points and the in-plane magnetic field strength in response to zero DC magnetic bias was visualized along the center plane of Ni sheet.



Phase field modeling: Details of the phase field modeling of the PZT/Ni ME composites are provided in the Supporting Information.32, 33 During the simulation process, the PZT layer is poled along its thickness direction, and an additional mismatch strain ε ij0 (r ) due to thermal expansion mismatch was introduced to the PZT layer in the phase field modeling. An in-plane magnetic field ∆H is applied along the Ni layer length direction, and the generated magnetostrictive strain elastically interacts with the poled piezoelectric layer, leading to a change in ferroelectric domain structure and polarization response ∆P. The ME coefficient is determined as αME = (1/ε0εr) ∆P/∆H.

3. RESULTS AND DISCUSSION

The X-ray diffraction (XRD) patterns of the as-deposited (AD) and laser annealed (LA) PZT films on Ni foil (Figure 2 (a)) showed peaks corresponding to a typical perovskite polycrystalline PZT film. The LA PZT film exhibited higher intensities for the XRD peaks as compared to the AD film, which can be attributed to the improved crystallinity of the PZT by laser annealing. Figure 2 (b) shows the Raman spectra obtained from the AD and LA PZT films. It can be observed that both the samples display similar peaks located at 210, 270, 560, 700 and 735 cm-1, which corresponds to the E(2TO), B1+E, A1(3TO), A1(3LO), and E(3LO) modes, respectively.34 These spectra represent the typical perovskite phase of PZT with composition close to the morphotropic phase boundary between PbZrO3 and PbTiO3. The higher intensity of Raman peaks of the LA PZT film further confirmed its greater degree of crystallinity as compared to the AD PZT film.35 The selected area electron diffraction (SAED) patterns collected from the AD and LA PZT films on Ni are shown in Figure 2 (c) and (d), respectively. The


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improvement in crystallinity of PZT by laser annealing is also indicated by the higher intensity of the diffraction rings in the SAED pattern, observed from corresponding radial profile, of the LA PZT film. During GSV deposition, the powder granules are sprayed onto the substrate, with high kinetic energy, in a vacuum chamber. Consequently, the fracturing and deformation of primary particles, due to high speed collision, results in a dispersed nanocrystalline phase in an amorphous matrix of the deposited film. The nanocrystalline regions in the AD PZT film may act as heterogeneous nucleation sites for the crystallization of amorphous phase under laser irradiation.28-31

Figure 2. (a) XRD patterns, (b) Raman spectra, and (c) and (d) SAED patterns of the asdeposited and laser annealed PZT films on Ni, respectively. Radial intensity profiles of SAED 9 Environment ACS Paragon Plus


patterns are also shown for both the samples (insets of (c) and (d)). All these structural characterization results indicate the improvement in crystallinity of the PZT film after laser irradiation.

The cross-sectional TEM image of the laser annealed PZT/Ni composite (Figure 3 (a)) indicates that the PZT film is well adhered to the Ni foil. Such a good interfacial bonding is favorable for efficient elastic strain transfer and enhanced ME coupling between PZT and Ni. The interfacial chemical homogeneity of LA PZT/Ni composite was investigated through the EDS analysis. The corresponding elemental distribution maps (Figure 3 (b) and (c)) confirm that there is no apparent diffusion or chemical reaction across the PZT/Ni interface after the laser annealing. The magnetization vs. magnetic field (M-H) behavior of the AD and LA PZT/Ni composites is shown in Figure 3(c). Both the samples exhibited almost identical but asymmetric M-H hysteresis loops (shifting of the hysteresis loop on field axis). The superposition of the anisotropic magnetization field and the applied AC magnetic field, during ME measurement, results in an internal magnetic bias in Ni. This gives rise to the self-biased ME coupling in the PZT/Ni composite. The EDS and M-H results suggest that the laser induced heating was localized to PZT layer and did not cause any damage to the Ni substrate. This clearly demonstrates the advantage of laser radiation for localized thermal treatment of piezoelectric films on heat-sensitive substrates.


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Figure 3. (a)-(c) Cross-sectional TEM micrograph and the corresponding EDS elemental mapping of the laser annealed PZT/Ni, respectively. (c) Magnetization hysteresis loops of the asdeposited and laser annealed PZT/Ni composites. No apparent diffusion or chemical reaction was observed across the PZT/Ni interface after laser irradiation, which is further supported by the identical magnetization behavior of the composite before and after laser annealing.

The dielectric properties of the AD and LA PZT films are compared in Figure 4 (a). Both the films exhibited low dielectric loss (tan δ) of 0.07-0.08 over a frequency range of 1 kHz to 1 MHz. The dielectric constant (εr) of the LA PZT film (~780 at 1 kHz) is much higher than that of the AD PZT film (~220 at 1 kHz). Figure 4 (b) shows the polarization-electric field (P-E) hysteresis loops of the AD and LA PZT films on Ni foil. The LA PZT film showed significantly better ferroelectric polarization than the AD PZT film. The measured remanent polarization (Pr) value of the LA PZT film was 40 µC/cm2, whereas the AD film exhibited a Pr value of 7



µC/cm2. This increased Pr after laser annealing indicates the enhanced piezoelectric response of the LA PZT film as piezoelectric constant is expressed as, dij = 2εQPr,where Q is the electrostriction constant.36 The higher dielectric and polarization properties of the LA PZT film than that of the AD film can be attributed to the improved crystallinity due to laser irradiation.31,37


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Figure 4. (a) Dielectric properties and (b) Polarization hysteresis loops of the as-deposited and laser annealed PZT films on Ni; (c) ME responses of the corresponding PZT/Ni composites. The electrical and ME properties of the PZT/Ni heterostructure were significantly improved due to the increase in crystallinity of the PZT film by laser irradiation.

The ME responses obtained from the AD and LA PZT/Ni composite samples are plotted in Figure 4 (c). The values of maximum αME of the AD and LA composites were measured to be 0.16 and 3.15 V/cm·Oe, respectively. The superior ME performance of the LA PZT/Ni can be mainly attributed to its improved electrical properties due to increased crystallinity of the PZT film. Furthermore, minimized mechanical damping at PZT/Ni interface, due to matching mechanical impedance of PZT (25-30 MRayl) and Ni (27 MRayl), can facilitate an efficient interfacial elastic strain transfer leading to better ME coupling in PZT/Ni composites.27, 38 The reported values of ME coefficient of PZT films deposited on Ni-based magnetostrictive substrates and their fabrication conditions are summarized in Table 1. From this table, it can be observed that the laser annealed PZT/Ni composite outperforms the furnace annealed (FA) and rapid thermal annealed (RTA) PZT films on Ni-based substrates. 21, 22, 27, 39 As can be seen from Figure 4 (c) and Table 1, the LA PZT films on Ni displayed fully selfbiased ME response, while, the FA and RTA PZT/Ni samples exhibited lower SME response than their respective maximum αME values. Because, the laser induced heat is localized to PZT layer only, the hysteretic magnetostrictive properties of the Ni substrate remained intact (Figure 3 (c)), leading to undiminished self-biased ME response, which is similar in both the asdeposited and laser-annealed PZT films on Ni. The absence of such a large SME response in FA



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and RTA PZT films on Ni, can be ascribed to the degraded magnetic properties of Ni at high annealing temperatures.21, 22 The lower demagnetization in the thin Ni foil substrate, used in this study, might also have played a significant role in achieving the enhanced ME coupling at zero bias.

Table 1. Comparision of reported ME voltage coefficients of PZT films deposited on Ni-based substrates and their annealing conditions. FA: furnance annealing, RTA: rapid thermal annealing, and LA: laser annealing.

ME composite

PZT/Ni thickness (µm)

Annealing condition

maximum self-biased αME αME % SME Reference (V/cm·Oe) (V/cm·Oe)

PZT/Pt/NiZnFe2O4 1/550

FA – 650 °C

0.14

0

Nil

[39]

PZT/Ni

1/200

RTA – 650 °C

0.22

0.2

90%

[21]

PZT/Pt/Ni

0.4/200

FA – 650 °C

0.772

0

Nil

[22]

PZT/LaNiO3/Ni

20/500

FA – 700 °C

1

0.2

20%

[27]

PZT/Ni

4/50

LA – 965 mW 3.15

3.15

100%

This work

The origin of internal magnetic bias generated in Ni was investigated by Zhou et al.12 using magnetic force microscopic analysis. It was revealed that the presence of macro-sized domains with long range ordering, in polycrystalline Ni, results in sizeable coercive field. Upon reorientation of the magnetic domains, higher field is required to switch back to the random state, leading to asymmetry in the magnetization (M) hysteresis of Ni, and further to an


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anisotropic magnetostriction and piezomagnetic coefficient, as ߣ௜௝ ∝ ‫ܯ‬ଶ and ‫ݍ‬௜௝ ∝

ௗெమ 39 . ௗு

While

the occurrence of self-biased ME coupling in Ni-based composites is related to the anisotropic λij of Ni and the corresponding remanent qij at Hdc = 0, the tunability of SME response has been found to depend on size-induced demagnetization and the resultant differential magnetic flux distribution in the magnetostrictive layer.12,13 A high flux concentration in the magnetic phase was observed to facilitate enhanced ME coupling under a low magnetic bias.40, 41 In ferromagnetic materials, the demagnetization field (Hd) is directly proportional to demagnetization factor (Nd) via Hd = MNd. Demagnetization in the magnetostrictive Ni layer depends mainly on its thickness, which is much smaller as compared to its transverse dimensions. By considering the influence of demagnetization field, the effective magnetic field (Heff) and the corresponding magnetic flux density (Beff) in the magnetic phase can be expressed using the following relations: 12, 13, 42 Heff = Hbias - Hd = Hbias - MNd

(1)

Beff = µ0(Heff +M) = µ0(Hbias +M) - µ0MNd

(2)

Here, µ0 is the permeability of free space. Figure 5(a) shows the variation of Nd as a function of Ni foil thickness (50-500 µm).43 It is clear that thinner Ni substrate exhibits smaller Nd, which could lead to lower Hd, and thus higher Heff. This implies that for achieving the same magnitude of Heff, one needs to apply smaller Hbias under lower Hd, and vice-versa. Accordingly, a maximum value of αME for a thinner Ni substrate can be obtained at a lower magnetic bias in comparison to its thicker counterpart.



Figure 5. (a) Variation of demagnetization factor with the thickness of Ni sheet; (b) FEM results of in-plane magnetic flux density distribution along the center plane of Ni sheets of varying thickness. Thinner Ni sheet (50 µm-thick) was found to exhibit higher magnetic flux concentration due to its lower demagnetization factor, compared to other thicker samples. The high flux density facilitates the realization of maximum αME under a low magnetic bias.

To verify the above correlation, the magnetic flux density distribution was simulated, using a finite element model (FEM),12, 13, 42 for Ni sheet with different thicknesses (50-500 µm) and the results are shown in Figure 5 (b). As seen from these results, the flux concentration in the Ni sheet significantly depends on its thickness. A thinner Ni sheet could display much stronger magnetic induction due to the lower Hd and higher Heff. Consequently, a larger value of αME was observed in the thinner Ni sheet under a lower Hbias. The shifting of the αME peak position towards the zero bias with decrease in thickness of the magnetostrictive layer was also observed in the case of Ni-based bulk ME laminates.12, 13 Based on the above findings, it can be inferred that the low Hd and high Heff in the thin Ni foil (50 µm-thick) could be responsible for achieving fully self-biased ME response in the PZT/Ni composite. In addition to the improved crystallinity


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of the PZT film, the strong magnetic induction in the thin Ni foil, could also facilitate the enhanced ME coupling in the composite. The larger differences between maximum αME and SME responses, reported in literature (Table 1) for PZT/Ni film composites with thicker Ni substrates, could be attributed to the higher demagnetization, and the resultant decrease in the magnetic flux concentration. Nevertheless, there seems to be a discrepancy between the SME responses of the PZT/Ni samples reported in literature,

22, 27

(Table 1), contrary to the trend

expected, with respect to the Ni layer thickness. As explained in the following section, through phase-field simulations of M-H and ME behavior of the PZT/Ni composite, besides the demagnetization in magnetostrictive layer, other factors such as mismatch strain (thermal and poling) in the laminate composite and its structural configuration (bilayer/multilayer, flexural/non-flexural moment) will also significantly influence the SME behavior. The differences in quantities of the above parameters of the PZT/Ni samples fabricated under different experimental conditions in other studies22, 27 (Table 1), might have led to the observed differences between their SME responses with respect to Ni layer thickness. In fact, it was observed in our phase-field simulations that at zero mismatch strain, the PZT/Ni composite with 100 µm thick Ni layer exhibited better SME response than the composite having 50 µm thick Ni layer (Figure 6 (a) and (b)), which is in contrary to the trend expected. Further, by considering some mismatch strain in the PZT/Ni composite with 50 µm thick Ni layer exhibited fully self-biased ME response in flexural state (Figure 6 (d). We believe that an optimized combination of magnetostrictive layer thickness (demagnetization/magnetic flux concentration), mismatch strain, and the laminate configuration will yield a highly enhanced selfbiased ME response in the composite.



Figure 6. Simulated domain structures (under zero magnetic bias), M-H hysteresis, and ME voltage coefficient α ME for laminate PZT/Ni composite under zero mismatch strain ε mis = 0 with varying demagnetization factors in Ni film plane: (a) N d = 0.008 , (b) N d = 0.014 , and (c)

N d = 0.026 , and under mismatch strain ε mis = 0.003 with the same demagnetization factor N d = 0.008 in (d) flexural and (e) non-flexural composite. Domain patterns are visualized by color maps with red, green, blue (RGB) components proportional to Px, Py, Pz in PZT layer and Mx, My, Mz in Ni layer, respectively.

To elucidate the underlying mechanisms responsible for the SME behavior, domain-level phase field modeling and computer simulation were employed. The phase field model for the ME composites developed in our previous work, which explicitly addresses the domain-level strain-


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mediated coupling between magnetization and polarization, was adopted.32, 33 The influence of demagnetization and mismatch strain on the shape of M-H loop and magnitude of remanent magnetization (Mr) and magnetic susceptibility (χr), which directly determines the SME behavior, were investigated. Results of the simulations indicate that the PZT/Ni heterostructure with 200 µm-thick Ni substrate (Nd = 0.026) show reduced magnetic hysteresis and lower Mr, thereby, displaying smaller αME at zero magnetic bias, as shown in Figure 6(c). Decreasing the Ni foil thickness to 100 µm (Nd = 0.014) was found to increase the magnetic hysteresis as well as the SME response (Figure 6(b)). However, further decrease in Ni foil thickness (50 µm) diminished the SME coupling due to the decreased magnetic susceptibility χ r (curve slope) at zero bias, in spite of an improved magnetic hysteresis and larger Mr (Figure 6(a)). These results indicate that there exists an optimal thickness of the Ni foil to realize the fully SME response in the PZT/Ni composite.

In the laminated PZT/Ni ME composite system, the total mismatch strain between the two layers can be determined by εmis = εth – εp, where εth is the thermal mismatch strain (resulting from different thermal expansion coefficients of the two layers and depends on the temperature change during the LA treatment), and εp is the poling mismatch strain (caused by the deformation of PZT layer due to the electrical poling). In an appropriate approximation, the planar mismatch strain of εmis=0.003 was considered in our modeling, which corresponds to an experimental deposition/annealing temperature of 600~700oC. Two representative laminate composites were considered in modeling of the mismatch strain effect: a flexural composite corresponding to bilayer PZT/Ni that could bend under the mismatch strain and a non-flexural composite corresponding to trilayer Ni/PZT/Ni (and multilayer structures) that could not bend. Figure 6(d) and (e) show the corresponding simulation results for these two types of composites with the



same demagnetization factor Nd = 0.008 (50µm-thick Ni). A significant difference was observed between the M-H loops as well as SME behaviors of the two composites, which resulted from different stress/strain distributions in the Ni layer.

For the non-flexural composite, a homogenous in-plane tensile stress is distributed within the Ni layer and stripe domain structures were formed in the Ni layer. The positive stressinduced perpendicular uniaxial anisotropy and the negative magnetostriction of Ni (λij = -40 ppm), significantly softened the M-H hysteresis and resulted in a nearly zero Mr, and thereby resulting in almost nil SME response (Figure 6(e)). For the flexural composite, both tensile (near the PZT/Ni interface) and compressive (far from the interface) stress are distributed along the length direction in the Ni layer due to bending, resulting in a mixing of in-plane (length direction) and perpendicular (thickness direction) magnetizations to produce non-zero magnetization remanence, and thus, larger SME response (Figure 6(d)). These above findings suggest the position and magnitude of maximum α ME and thereby the SME response could be tailored by tuning demagnetization factor, mismatch strain, and flexural moment in the PZT/Ni heterostructure. An optimum combination of these factors should be considered to realize a fully self-biased ME response in the composite. It is interesting to see that both the AD and LA PZT/Ni samples displayed fully self-biased ME coupling, which implies that the crystallinity of PZT film or thermal stress mismatch do not critically affect the SME behavior of the PZT/Ni heterostructure.

4. CONCLUSION In summary, a magnetoelectric heterostructure of PZT/Ni was fabricated by depositing PZT thick film on thin flexible Ni foil. Deposition of PZT film at room temperature using


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granule spray in vacuum process combined with localized annealing of the film through laser radiation provided feasible fabrication approach towards overcoming the issues related to the high temperature processing of ceramic films on metal substrate. The laser induced crystallization of PZT film significantly contributed to the improvement of dielectric, ferroelectric and ME properties of the PZT/Ni heterostructure. The laser annealed PZT film on Ni exhibited a highly enhanced and fully self-biased ME coupling of 3.15 V/cm·Oe, which is the highest among all the reported values in literature for similar systems. The demagnetization in the Ni foil, mismatch strain, and flexural moment in the PZT/Ni bilayer composite were found to greatly influence its self-biased ME response.

ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publication website at DOI: Phase field modeling of laminate PZT/Ni ME composites

AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected] *E-mail: [email protected]

Notes The authors declare no competing financial interest.



ACKNOWLEDGEMENTS This research work was supported by the National Research Foundation of Korea (Grant No. NRF-2016R1A2B4011663); Korea Institute of Materials Science (KIMS) internal R&D program (Grant No. PNK5061); and the U.S. Office of Naval Research Global (Grant No. N62909-16-12135). D.M. would like to acknowledge support from Office of Basic Energy Science, Department of Energy (Grant No. DE-FG02-06ER46290). S.P. would like to acknowledge support from Office of Naval Research (Grant No. N00014-16-1-3043). Y.S.O. was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT & Future Planning (NRF-2015R1C1A1A01055964 and NRF-2016K1A3A7A09005338).

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Enhanced Self-Biased Magnetoelectric Coupling in Laser-Annealed

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