Voltage-Impulse-Induced Non-Volatile Ferroelastic Switching of Ferromagnetic Resonance for Reconfigurable Magnetoelectric Microwave Devices

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1 Ming Liu, * Brandon M. Howe, Lawrence Grazulis, Krishnamurthy Mahalingam, Tianxiang Nan, Nian X. Sun, and Gail J. Brown Voltage-Impulse-Induced Non-Volatile Ferroelastic Switching of Ferromagnetic Resonance for Reconfigurable Magnetoelectric Microwave Devices The central challenge in tunable magnetic microwave devices lies in finding an energy efficient way to perform wide ferromagnetic resonance (FMR) tuning in a reversible and reproducible manner using a control voltage, rather than a current-driven electromagnet. [ 1 7 ] This requires the creation of novel materials and functionalities, while integrating into non-volatile electronic devices. [ 8 12 ] Multiferroic heterostructures, [ 13, 7, ] exhibiting strong strain-mediated magnetoelectric (ME) coupling between distinct ferromagnetic and ferroelectric phases, have shown great promise for frequency agile microwave applications. [ 25, 1, 3 4, ] In these materials, a single control parameter of voltage-induced piezo-strain, arising from ferroelectrics, is used to in situ introduce an effective magnetic anisotropy [ 26, ] and subsequent FMR frequency shift, in elastically-coupled ferromagnetic phases via the magnetoelastic effect. [ 6, 33, ] Therefore, devices based upon such materials are, in principle, light-weight, fast, and energy efficient, therefore overcoming some of the intrinsic limitations in conventional microwave components in addition to providing new functionality. In most prototype ME microwave devices such as tunable filters, [ 43 ] resonators, [ 39 ] and phase shifters, [ 29 ] tuning of FMR frequency has been achieved through the use of a linear piezo response, whereby the strain is directly proportional to the applied voltage, without electric polarization switching. [ 44 ] Upon removing the electric field, the FMR decays to the initial state. [ 27, 44 ] While these devices point to a pathway for enhancing FMR tunability, reversible and non-volatile tuning of FMR using electric fields has thus far remained relatively unexplored. However, non-volatile tuning is indispensable from a device application point of view. In non-volatile switching, the FMR states remain in a stable remnant state after the controlled impulse voltage has been switched off. To achieve this goal, we need to focus on changing FMR by using non-linear latticestrain effects arising from the ferroelastic domain switching Dr. M. Liu, Dr. B. M. Howe, L. Grazulis, Dr. K. Mahalingam, Dr. G. J. Brown Materials and Manufacturing Directorate Air Force Research Laboratory Wright-Patterson Air Force Base, OH , USA ming.liu.ctr@wpafb.af.mil T. Nan, Dr. N. X. Sun Electrical and Computer Engineering Department Northeastern University Boston, MA 02115, USA and selectively control of the switching pathway in a reversible and reproducible manner. Recently, Zhang et al. reported non-volatile magnetization tuning using ferroelastic domain switching in (001) oriented PMN-PT. [ 45 ] However, due to the 71 and 180 polarization rotation-induced equivalent strain states in two stable polarization directions, only 109 ferroelastic switching contributes to the in-plane lattice strain, and merely covers 26% of the entire poled area. Such small coverage may cause inhomogeneity of in-plane deformation, leading to a lack of switching control due to the complicated competition among those polarization rotations. In this work, we used an unique ferroelastic switching pathway in (011)-oriented PMN-PT (0.71Pb(Mg 1/3 Nb 2/3 ) O PbTiO3 ) single crystals that allows up to 90% of the polarization to rotate from an out-of-plane to a purely in-plane direction (71 and 109 switching), thereby producing two distinct, stable, and electrically-reversible lattice strain states. Voltage-impulse switching between these remnant stain states is demonstated and results in a highly energy-efficient, nonvolatile tuning of FMR frequency up to 2.3 GHz in elasticallycoupled amorphous FeCoB films deposited on PMN-PT (011). Domain distortion, polarization switching pathway, and the resultant lattice strain produced in response to in situ verticallyapplied voltages in PMN-PT (011) are clearly presented using reciprocal space mapping (RSM) and piezoforce microscopy (PFM) measurements. Up to 90% of polarization is elastically switched in a stable and reversible manner (109 and 71 ferroelastic switching), which contributes to a strong homogeneous lattice strain through the entire sample. These results provide a framework for realizing reconfigurable, compactable, light-weight, and ultra-low power non-volatile electronics and microwave devices being deployed in aircraft, satellite and portable communication systems. Multilayer films with the structure of Au(5 nm)/ Fe 60 Co20B 20 (50 nm)/ti(5 nm) were deposited on (011)-oriented single crystalline PMN-PT substrates using magnetron sputtering. FeCoB alloy is selected due to its high magnetostriction constant ( 60 ppm) and small FMR linewidth in the amorphous phase. [ 46 ] The amorphous nature of FeCoB, as well as a strong mechanical coupling between films and substrates is confirmed by high-resolution TEM images as shown in Figure S1 in Support Information. An in-plane magnetic anisotropy is found in as-grown films ( Figure 1 a), where the hysteresis loop obtained along the [0 11] direction is nearly square, while the [100] direction shows hard axis behavior. The anisotropy field is determined to be 16 Oe and could arise due to the anisotropic 1

2 Figure 1. (a) In-plane magnetic hysteresis loops of FeCoB/PMN-PT (011). Insets are schematic (upper left) and FMR spectra (bottom right). (b) Schematic of FMR measurement for (c-f). The sample is laid face down on an S-shape co-planar waveguide. Magnetic fi elds are applied in the [100] direction and electric fi elds are applied along the [011] direction. (c) Electric fi eld dependence of the FMR frequency in fi eld sweeping mode. (d) Electric fi eld dependence of the FMR fi eld in frequency sweeping mode. (e) FMR frequency responses under unipolar (red) and bipolar (blue) sweeping of electric fi elds at room temperature. (f) Voltage-impulse-induced non-volatile switching of FMR frequency. residual stress caused by the difference in lattice parameters along the in-plane [0 11] and [100] directions. The magnetization is determined to be 1230 emu cm 3. Electric-field-induced FMR modulations in FeCoB/PMN-PT(011) were demonstrated using a custom made coplanar-waveguide (CPW) FMR test unit. Figure 1 b shows the measurement schematic. In layered magnetoelectric heterostructures such as FeCoB/PMN-PT (011), an electric-field-induced effective magnetic field, H eff is expressed as H eff = 3λ(σ 100 σ 0 11) [ (1), 44 ] where σ M 100 and σ 0 11 s are the in-plane piezo-stress, λ is the magnetostriction constant of FeCoB films, and M s is the saturation magnetization. This field is able to manipulate FMR as described in the Kittel equation f = γ H r + H eff H r + H eff + 4π M s (2), ) [ 2 ] where γ is the gyromagnetic ratio 2.8 MHz Oe 1. Details are described in ref. [ 44 ]. Figure 1 c,d shows the electric-field dependence of the FMR spectra in field sweeping mode at 9.3 GHz and frequency sweeping mode at a bias field of 720 Oe. Magnetic fields are applied along the in-plane [100] direction (See Figure S2 in Support Information for the case that the magnetic field is applied along the [0 11] direction). Upon applying an electric 2

3 field of 6 kv cm 1 on a negatively-poled FeCoB/PMN-PT (011) sample, a slight increase in resonance field from 665 Oe to 715 Oe (Figure 1 c) and decrease in frequency from 9.9GHz to 9.45 GHz (Figure 1 d) is observed, with the linewidth remaining the same. This is due to the linear piezo effect that leads to a linear polarization elongation as well as an in-plane compressive strain along the [100] direction. Note the positive magnetostriction of FeCoB, a negative effective magnetic anisotropy H eff is produced and leads to the increase in the resonance field and reduction in the frequency as described in Equation 2. [ 2, 44 ] However, as a small positive E-field of 1.5 kv cm 1 is applied on a negatively-poled sample, the tunable range of FMR increases dramatically by Δ H r = 320 0e in field sweeping mode and Δ f = 2.3 GHz in frequency sweeping mode (Figure 1 c,d). We relate this enhanced FMR tunability to the 71 and 109 ferroelastic domain switching, whereby the polarization is suppressed from the out-of-plane direction to the in-plane direction. This polarization rotation results in a strong ferroelastic tensile strain along the in-plane [0 11] direction. Figure 1 e shows the resonance frequency in response to in situ electric fields applied normal to the sample. A Butter-fly curve (blue) is observed as cycling triangle electric fields within 6 kv cm 1, showing a maximum tunable frequency range of 2.3GHz. Upon applying a positive electric field on a negatively-poled FeCoB/PMN-PT (011), a giant frequency jump takes place near the coercive field of 1.5 kv cm 1. In this process, the polarization undergoes 71 and 109 ferroelastic switching from the out-ofplane to the in-plane direction, and is accompanied by an in-plane lattice strain induced by the domain distortion. Therefore, the hysteresis loop of the FMR frequency as a function of the electric field is observed (Figure 1 e). Similar to magnetic memory, two stable and reversible frequency remnant states A and B would facilitate the realization of non-volatile frequency switching by reversing the applied electric field at the coercive field. Figure 1 f shows voltage-impulse-induced non-volatile FMR frequency tuning in FeCoB/PMN-PT (011). As the PMN-PT (011) is subjected to an impulse of 6 kv cm 1, the remnant strain in poling state A is retained and results in a maximum FMR frequency of 9.9 GHz. Upon applying an impulse field of 1.5 kv cm 1, the resonance frequency is reduced to 7.6 GHz, indicating that the strain state is switched to remnant state B. Furthermore, reversing the bias again to an electric impulse of 6 kv cm 1 returns the FMR frequency to its maximum value. Besides non-volatile switching the FMR frequency between the maximum and minimum values, any frequency between them is reachable by choosing a proper positive electric impulse. For example, the non-volatile frequency switching between 10 GHz and 8.5 GHz can be realized by applying pulsed electric fields of 6 kv cm 1 and 1.4 kv cm 1 as shown in Figure 1 (f) (green). In order to confirm that non-volatile tuning of FMR in FeCoB films is attributed to the stable and reversible polarization switching in PMN-PT (011), high resolution x-ray diffraction (HRXRD) measurements were used to determine the polarization switching pathway and lattice strain in response to in situ electric fields. Figure 2 shows the electric field dependence of the reciprocal space maps (RSMs) in the vicinity of the (022) and (002) reflections of the PMN-PT (011) substrate. Due to the rhombohedral symmetry of the single crystal PMN-PT used in this work, [ 47 ] there are eight polarization directions pointing along the body diagonals of the pseudocubic unit cell with four structural domains (r1,r2,r3,r4), as shown in a schematic in Figure 2 a. Under vertically-applied electric fields, the polarization switching pathways include 71 ferroelastic switching from r1 + to r3 + /r4, 109 ferroelastic switching from r1 + to r4 + /r3 and 180 ferroelectric switching from r1 + to r1. According to the lattice parameter and rhombohedral distortion, [ 47 ] structural domains r1/r2 (polarization points to the out-of-plane direction) and r3/r4 (polarization stays in the plane) can be distinguished by the spot distribution about the PMN-PT (022) reflection, due to the difference in d -spacing of these distorted domains. For the unpoled state of PMN-PT (011) (the first column in Figure 2 ), a single broad spot is observed in both (022) and (002) reflections (Figure 2 e and 2 i). Analysis of these RSMs suggests that two possible domain structures, r3 and r4, are dominant in the unpoled state, and most of the polarization lies in the plane. As the sample is vertically poled with a strong positive voltage, the RSM in Figure 2 f shows an addition high intensity (022) reflection spot with a lower Q 022 value, corresponding to the r1/r2 domain structures. Meanwhile, the intensity of the spot corresponding to r3/r4 reduces dramatically. This indicates that 71 and 109 ferroelastic polarization switching from the in-plane direction to the out-of-plane direction takes place and results in a large out-of-plane lattice strain. After a small negative electric field of 1.5 kv cm 1 is applied and removed, the domain distortion returns to r3/r4 and polarization is suppressed from the out-plane direction to the in-plane direction (Figure 2 g). As a large positive electric field of 5 kv cm 1 is applied and then switched off, the domain structure is again switched back to r1/r2 (Figure 2 h). Therefore, a stable and reversible ferroelastic domain switching pathway is confirmed, which enables polarization rotation between the in-plane direction and the out-ofplane direction. In comparison to (001)-oriented PMN-PT, where only 109 domain switching facilitates lattice strain and merely covers 26% of the poled area, (011)-oriented PMN-PT allows both 71 and 109 ferroelastic switching, contributing to large in-plane strain and covering up to 90% of the entire poled area (from peak intensity analysis). We note that all the (002) reflections (Figure 2 i- 2 j) show a single peak with a same Q 002 value due to the same d -spacing of (002) planes for all possible domain distortions. To quantify the in-plane lattice strain in response to domain switching, the lattice parameters along different orientations under various vertically-applied electric fields were measured. Figure 3 a shows the schematics of polarization switching from the out-of-plane direction to the in-plane direction upon applying an electric field, and its induced changes in the (022) diffraction peak position. Before switching (Configuration I), all polar vectors point upward, corresponding to the r1/r2 domain structures. A single (022) diffraction peak is observed with the out-of-plane lattice parameter c = Å. As a negative electric field is gradually applied to the positively-poled PMN-PT (Configuration II), polarization vectors partially rotate from the out-of-plane direction to the in-plane direction, and all possible domain distortions exist. Therefore, two distinct diffraction peaks with c- axis lattices of c1 = Å and c2 = Å are obtained, corresponding to the r1/r2 and r3/r4 domains, respectively. Please note that the effective lattice constant during domain switching is determined by the two diffraction peaks 3

4 Figure 2. Schematics of domain structures and reciprocal space maps (RSMs) about (022) and (002) reflections of PMN-PT(011) under various applied electric fi elds and thus poling states. The fi rst column (a,e,i) is for the unpoled state. The second column (b,f,j) is for the positive poling state with up to 90% of polarization pointing upward. The third column (c,g,k) is after applying an negative electric fi eld of 1.5 kv cm 1 and then switching it off. The fourth column (d,h,l) is achieved by applying a positive electric fi eld of 5 kv cm 1 and then switching it off. intensities as well as their lattice parameters. After the polarization is completely suppressed to the in-plane (Configuration III), the domains remain at r3/r4 and present a single (022) diffraction peak with c -axis lattice of Å (See Figure S3 in Support Information for details). Figure 3 b,c shows electric field dependence of lattice parameters along the [001] ( a-axis) and [011] ( c -axis) directions. A slight a -axis lattice change of less than 0.05% is observed, revealing that the lattice strain along [001] arising from domain switching is negligible. However, a large c -axis lattice change by 0.25% is found as the domain switched between r1/r2 and r3/r4 (Figure 3 c). These two remnant strain states A and B, corresponding to r1/r2 and r3/r4 respectively, are stable and switchable. Figure 3 d shows the inplane strain curve along [0 11] direction deduced from the a- and c -axis lattice changes. A large tensile strain of 0.22% along the in-plane [0 11] direction is produced, while a negligible lattice change takes place along [100] direction. The in-plane strain in response to the ferroelastic domain switching is consistent with the observation of the FMR as a function of electric field as shown in Figure 1. This highlights the extraordinary nature of poling PMN-PT (011) in comparison to PMN-PT (001), [ 45 ] with dominant ferroelastic switching in PMN-PT (011) occurring due to the existence of stable in-plane polarization states. To further confirm the observed domain switching dynamics at the surface of (011)-oriented PMN-PT, we imaged the polarization domains using PFM. Figure 4 shows the vertical PFM (VPFM) phase and topography images in response to various vertically-applied voltages. For the unpoled state at room temperature, the polarization vectors on the surface randomly point to the eight body diagonals of the pseudocubic cell, and is confirmed by the VPFM phase images (Figure 4 a-d) and lateral PFM (LPFM) images (Figure S4). When a poling voltage of + 15 V is applied on the tip, both ferroelectric switching (180 ) and ferroelastic switching (71 and 109 ) take place and all the polar vectors point upward in the poled area (blue box). As a negative voltage of 5 V is subsequently applied, the polarization vectors change to an intermediate state with a large majority of polarization lying purely in the (011) plane, as shown in Figure 4 c. As the voltage is further increased to 20 V, all of the polarizations point downward, as shown in Figure 4 d. These results indicate that upon gradually increasing the poling bias, a two-step 71 or 109 ferroelastic switching takes place, which first makes the polarization components completely rotate from the upward direction into (011) in-plane directions, and subsequently points to the downward direction at the end of 4

5 Figure 3. a) Schematics of ferroelastic domain switching from r1/r2 to r3/r4 (upper left) and the evolution of the theta-2theta scans along the (022) peak (upper right) during the domain switching. The lattice parameters along the [001], [0 11] and [011] directions are marked as a, b, and c, respectively. b,c,d) are the lattice parameters and strains along the [001],[011] and [0 11] directions as a function of the applied electric fi elds. the reversal. Figure 4 e- 4 h are surface topography images of the same scan areas. Distinct contrast changes in switching areas (blue box) under different poling states indicates that large displacements are produced during ferroelastic domain switching. The relative height changes in switched domains are plotted in line profiles, as shown in Figure 4 i- 4 l. For the poled states, where the polarization points to the out-of-plane direction, the displacement in the [011] direction is enhanced (Figure 4 j and 4 l). However, for the unpoled state and the state with most of the polarization aligned in-plane (Figure 4 i 5

6 Figure 4. Probe-bias-dependent switching dynamics in (011) oriented single crystal PMN-PT. (a,b,c,d) are the out-of-plane vertical PFM (VPFM) phase images upon applying different voltages to the square area outlined by a blue dashed line. (e,f,g,h) are corresponding topography images. (i,j,k,l) are topography line profi les across the switching area for each poling state. The height change indicates an expansion or contraction in the out-of-plane direction. and 4 k), the height shows slightly differences. These results reveal a large deformation during the non-volatile polarization switching between the in-plane direction and the out-of-plane direction. In conclusion, we have demonstrated voltage-impulseinduced non-volatile FMR modulation in FeCoB/PMN-PT (011), showing a remarkable FMR tunable range and the capability to freely tune the resonance frequency to the required position. This is achieved using a unique ferroelastic polarization switching pathway in PMN-PT (011) to maximize ferroelastic strain and realize two distinct, stable, and completely reversible lattice strain states. These results point to opportunities for electrical tuning of strain-sensitive properties in all materials, while providing a framework for realizing reconfigurable, frequency-agile, non-volatile, and highly energy-efficient tunable electronics and microwave devices. Experimental Section Multilayer films with the structure of Au(5 nm)/fe 60 Co 20 B 20 (50 nm)/ Ti(5 nm) were deposited on (011) oriented single crystalline PMN-PT substrates [7.5 mm(l) 7.5 mm(w) 0.5 mm(t)] at room temperature using magnetron sputtering. High resolution transmission electron microscopy (HRTEM) was used to characterize the heterostructure. Domain structure as well as switching pathway in PMN-PT (011) was characterized by reciprocal space mapping (RSM) using a triple axis high-resolution x-ray diffractometer. Domain switching dynamics at the surface of PMN-PT (011) were characterized by piezoforce microscopy (PFM). The magnetic hysteresis loops were measured using a superconducting quantum interference device (SQUID) magnetometer at room temperature. The ferromagnetic resonance spectra were measured using our custom made wide-band and narrow-band FMR test unit. Acknowledgements This work was supported by the AFRL Materials & Manufacturing Directorate Applied Metamaterials Program. Received: May 2, 2013 Revised: May 30, 2013 Published online: [1 ] G. Srinivasan, Ann. Rev. Mater. Res. 2010, 40, 153. [2 ] M. Liu, O. Obi, J. Lou, Y. J. Chen, Z. H. Cai, S. Stoute, M. Espanol, M. Lew, X. Situ, K. S. Ziemer, V. G. Harris, N. X. Sun, Adv. Funct. Mater. 2009, 19, [3 ] Ü. Özgür, Y. Alivov, H. Morkoç, J. Mater. Sci.: Mater. Electron 2009, 20, 789. [4 ] M. Liu, Z. Zhou, T. Nan, B. M. Howe, G. J. Brown, N. X. Sun, Adv. Mater. 2013, 25, [5 ] A. M. Hermann, R. M. Yandrofski, J. F. Scott, A. Naziripour, D. Galt, J. Price, J. Cuchario, R. K. Ahrenkiel, J. Supercond. 1994, 7,

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