This is the pre-peer reviewed version of the following article: Zhao W., Kim J., Huang X., Zhang L., Pesquera D., Velarde G.A.P., Gosavi T., Lin C.-C., Nikonov D.E., Li H., Young I.A., Ramesh R., Martin L.W.. Low-Voltage Magnetoelectric Coupling in Fe0.5Rh0.5/0.68PbMg1/3Nb2/3O3-0.32PbTiO3 Thin-Film Heterostructures. Advanced Functional Materials, (2021). . : - . 10.1002/adfm.202105068, which has been published in final form at https://dx.doi.org/10.1002/adfm.202105068. This article may be used for non-commercial purposes in accordance with Wiley Terms and Conditions for Use of Self-Archived Versions. 1 Low-Voltage Magnetoelectric Coupling in Fe0.5Rh0.5/ 2 0.68PbMg1/3Nb2/3O3-0.32PbTiO3 Thin-Film Heterostructures 3 4 Wenbo Zhao1*, Jieun Kim1*, Xiaoxi Huang1, Lei Zhang1, David Pesquera1,4, Gabriel A. P. Velarde1, 5 Tanay Gosavi3, Chia-Ching Lin3, Dmitri E. Nikonov3, Ian A. Young3, Ramamoorthy Ramesh1,2, Lane 6 W. Martin1,2,† 7 1 8 Department of Materials Science and Engineering, University of California, Berkeley, CA, USA 2 9 Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, CA, USA 3 10 Components Research, Intel Corporation, Hillsboro, OR, USA 4 11 Catalan Institute of Nanoscience and Nanotechnology (ICN2), CSIC and BIST, Campus UAB, 12 Bellaterra, 08193 Barcelona, Spain 13 14 * These authors contributed equally to the work; † lwmartin@berkeley.edu 15 16 ABSTRACT 17 The rapid development of computing applications demands novel low-energy consumption 18 devices for information processing. Among various candidates, magnetoelectric heterostructures 19 hold promise for meeting the required voltage and power goals. Here, we demonstrate a route to 20 low-voltage control of magnetism in 30 nm Fe0.5Rh0.5/100 nm 0.68PbMg1/3Nb2/3O3-0.32PbTiO3 21 (PMN-PT) heterostructures wherein the magnetoelectric coupling is achieved via strain-induced 22 changes in the Fe0.5Rh0.5 mediated by voltages applied to the PMN-PT. We describe approaches to 23 achieve high-quality, epitaxial growth of Fe0.5Rh0.5 on the PMN-PT films and, a methodology to 24 probe and quantify magnetoelectric coupling in small thin-film devices via anomalous Hall effect 1 1 studies. By comparing the spin-flop field change induced by temperature and external voltage, the 2 magnetoelectric coupling coefficient is estimated to reach ~7.08 x 10-8 s/m at 325 K while applying 3 a -0.75 V bias. 4 KEYWORDS: Multiferroic heterostructures; Piezo-strain effect; Magnetoelectric coupling; Non- 5 volatile; anomalous hall effect. 6 INTRODUCTION 7 Moore’s Law1 and Dennard scaling2 have been important guiding principles in the semiconductor 8 industry. With the continued decrease of critical dimensions in logic and memory devices, however, 9 there has been an increasing challenge to meet the requirements of these guiding principles, raising 10 important challenges for the future of integrated circuit (IC) design.3 The strong desire to identify 11 beyond-CMOS devices and technologies4–6 has driven increasing attention to a range of alternative 12 computing platforms, including systems based on electron,7–9 spin,10,11 ferroelectric,12 strain,13 and 13 other phenomena to overcome the “Boltzmann Tyranny” in electronic systems. Spin-based logic 14 has emerged as one of the leading options4,6 due to an intriguing combination of non-volatility, 15 higher logic efficiency,14 and the potential for built-in memory-in-logic and/or logic-in-memory 16 function. Traditional spin-logic devices, however, have limitations – including large energy 17 dissipation during current-driven magnetic field or spin-transfer torque generation and the fast spin 18 decoherence of many semiconductors results from spin-orbit coupling.15 19 Recently, a magnetoelectric spin-orbit (MESO) logic-in-memory device has been proposed 20 to address these challenges.16 Taking advantage of the fact that a charge input/output is easier to 21 accomplish, the input signal in the MESO device is a voltage input that is converted by the 2 1 magnetoelectric into a spin signal. After the logic operations are completed in the spin domain, the 2 spin is then back-converted to a charge signal using the inverse spin Hall effect.16 For the 3 magnetoelectric module, three mechanisms have been proposed to accomplish the required 4 function: strain, exchange-coupling, or charge-induced effects.17–24 Each of these presents 5 advantages and drawbacks; for example, exchange-bias coupling of the magnetoelectric to a 6 ferromagnet implicitly breaks time-reversal symmetry but is an interface mediated coupling 7 mechanism which means that the strength of the coupling is delicate. In contrast, strain-mediated 8 approaches are longer range, but have the disadvantage of not breaking time symmetry and hence 9 are useful in manipulating the magnetic anisotropy. Compared with the two other approaches, 10 strain in composite heterostructures provides the potential for robust coupling in large area and 11 thickness devices,17,21,25 and has made this approach one of the more attractive options. Numerous 12 studies have explored the potential for magnetoelectric coupling in ferromagnetic/piezoelectric 13 heterostructures including Co/0.7PbMg1/3Nb2/3O3-0.3PbTiO3,26,27 Co0.4Fe0.4B0.2/ 14 0.7PbMg1/3Nb2/3O3-0.3PbTiO3,18,28 Fe0.5Rh0.5/0.72PbMg1/3Nb2/3O3-0.28PbTiO3,21 etc. These 15 works, however, made use of bulk or thick-film piezoelectrics which require relatively high 16 voltages (typically >250 V, corresponding to 5 kV/cm for 0.5-mm-thick piezoelectrics) and, thus 17 large energies, to actuate and thus do not meet the growing call for low-power, energy-efficient 18 devices.29 19 Here, magnetoelectric coupling is explored in epitaxial multiferroic heterostructures of 30 20 nm Fe0.5Rh0.5 / 100 nm 0.68PbMg1/3Nb2/3O3-0.32PbTiO3 (PMN-PT) and 30 nm Fe0.5Rh0.5 / 10 nm 21 MgO / 100 nm PMN-PT / 25 nm Ba0.5Sr0.5RuO3 (BSRO) / NdScO3 (NSO) (110) heterostructures. 3 1 The thin MgO-buffer layer was employed to protect the PMN-PT thin films during high- 2 temperature Fe0.5Rh0.5 deposition in reducing conditions. It is found that in both cases, high-quality, 3 epitaxial, (001)-oriented films of Fe0.5Rh0.5 can be grown on the PMN-PT thin films. Subsequent 4 temperature-dependent magnetization and resistivity studies reveal the ability to produce near- 5 room-temperature phase transitions in the Fe0.5Rh0.5. A methodology to probe and quantify the 6 magnetoelectric coupling in small thin-film devices via anomalous Hall effect studies is presented 7 wherein one can compare the spin-flop field change induced by temperature and external voltage 8 to extract a measure of the magnetoelectric coupling coefficient. Values approaching ~7.08 x 10-8 9 s/m at 325 K while applying just -0.75 V bias are achieved. The results demonstrate the potential 10 for near room temperature low-voltage control of magnetism in Fe0.5Rh0.5 via piezoelectric strain 11 in PMN-PT thin films. Future directions and possible improvements to these devices are also 12 discussed. 13 14 RESULTS AND DISCUSSION 15 Here, we focus on using Fe0.5Rh0.5 to demonstrate the potential for strain-mediated magnetoelectric 16 coupling. This material has received considerable attention as a potential candidate for electrical- 17 field control of magnetization due to its metamagnetic transition from an antiferromagnetic (AFM) 18 phase to a ferromagnetic (FM) phase around T* ≈ 350 K.30–32 Previous studies demonstrated that 19 epitaxial Fe0.5Rh0.5 thin films can be grown on single-crystal substates of piezoelectrics such as 20 BaTiO3 and PMN-PT and that field-induced strains from the substrates were able to shift the AFM- 21 to-FM phase transition and induce large electroresistance effects.17,21 We explored the optimization 4 1 of Fe0.5Rh0.5 growth in three heterostructure variants: 2 nm Pt / 30 nm Fe0.5Rh0.5 / MgO (001) 2 (henceforth, Fe0.5Rh0.5/MgO; Fig. 1a), 2 nm Pt / 30 nm Fe0.5Rh0.5 / 100 nm PMN-PT / 25 nm 3 Ba0.5Sr0.5RuO3 / NdScO3 (110) (henceforth, Fe0.5Rh0.5/PMN-PT; Fig. 1b), 2 nm Pt / 30 nm 4 Fe0.5Rh0.5 / 10 nm MgO / 100 nm PMN-PT / 25 nm Ba0.5Sr0.5RuO3 / NdScO3 (110) (henceforth, 5 Fe0.5Rh0.5/MgO/PMN-PT; Fig. 1c). Growth on MgO (001) substrates was used for internal 6 reference due to the excellent lattice matching between Fe0.5Rh0.5 (body-centered cubic, a = 0.423 7 nm) and MgO (rock-salt structure, a = 0.421 nm) and prior works showing that high-quality films 8 can be produced therein.17,21 Details of the growth optimization are provided (Experimental 9 Section and Suppl. Fig. S1). The PMN-PT films were produced following conditions established 10 in previous studies using pulsed-laser deposition.33,34 The insertion of the MgO-buffer layer was 11 explored in case the growth of the Fe0.5Rh0.5 adversely effected the properties of the PMN-PT. 12 Multiple thicknesses of MgO-buffer layer were explored (3-10 nm; data for growths with 3 nm 13 MgO are also provided, Suppl. Fig. S2), and it was found that the slightly thicker MgO layers 14 provided for a better template layer upon which to produce high-quality Fe0.5Rh0.5. Following 15 growth, the films were transferred to a UHV sputter deposition system for the platinum and 16 Fe0.5Rh0.5 deposition at high temperatures. In all cases, it was found that high-quality, epitaxial 17 (00l)-oriented Fe0.5Rh0.5 films could be produced with the best results corresponding to growth at 18 600°C (Experimental Section, Fig. 1d). Estimates of the Fe0.5Rh0.5 quality and ordering can be 19 obtained by comparing the diffraction intensity for the 001- and 002- diffraction conditions,35 and 20 via this analysis the Fe0.5Rh0.5/PMN-PT heterostructures were found to be as good or better than 21 the Fe0.5Rh0.5/MgO heterostructures (Suppl. Fig. S3). Subsequent studies of film stoichiometry via 5 1 Rutherford backscattering spectrometry (Experimental Section) indicate that the films grown at 2 the optimized conditions have chemistry Fe0.49Rh0.51 (Fig. 1e), within error, the same as the source 3 target and provided an expected AFM-to-FM phase transition at ~350K.30,31 4 Having established the ability to create high-quality, epitaxial Fe0.5Rh0.5/PMN-PT 5 heterostructures, we proceeded to explore the magnetic and transport properties as a function of 6 temperature to further gauge the quality of the magnetic layer. Magnetization and resistivity (ρ) 7 studies as a function of temperature show the expected first-order phase transition behavior for all 8 heterostructures (Fig. 2), consistent with previous reports.17,21 For the Fe0.5Rh0.5/MgO 9 heterostructures, the AFM-to-FM phase transition occurs from 280-380 K (mid-point ~330 K) with 10 a hysteresis that is ~40 K wide at the mid-point (Fig. 2a,b). At temperatures below ~250 K, the 11 magnetization is essentially zero, indicating that the Fe0.5Rh0.5 has been fully transformed to the 12 AFM phase. This is accompanied by a change in the resistivity (Fig. 2b). These Fe0.5Rh0.5/MgO 13 heterostructures compare favorably to similar structures reported in the literature and suggest we 14 have high quality Fe0.5Rh0.5 under the optimized growth conditions (Suppl. Fig. S4). For the 15 Fe0.5Rh0.5/PMN-PT heterostructures, the larger lattice mismatch between the Fe0.5Rh0.5 and 16 potential instability of the PMN-PT to high-temperature growth in reducing environments is found 17 to be detrimental for the manifestation of idealized magnetic and transport properties (Fig. 2c,d). 18 Although no impurity phases were detected from X-ray diffraction, the magnetization studies 19 reveal an additional ferromagnetic phase contribution, as indicated by the non-zero magnetization 20 value observed at low temperatures (< 250 K), the wide temperature range of phase transition, and 21 an enhanced mid-point temperature for the transition (Fig. 2c). This further suppresses the 6 1 difference between the resistivity of the two states (Fig. 2d). Finally, for the Fe0.5Rh0.5/MgO/PMN- 2 PT heterostructures, there is a (partial) improvement of the properties compared to the Fe0.5Rh0.5- 3 /PMN-PT heterostructures (Fig. 2e,f). While the width of the phase transition (hysteresis) is still 4 quite diffuse and the mid-point temperature for the transition is enhanced, the magnetization does 5 reach zero at low temperatures; suggesting an improved crystalline quality for the Fe0.5Rh0.5 upon 6 inclusion of the MgO-buffer layer. Nonetheless, the resistivity behavior is less clean than for films 7 grown directly on MgO substrates. 8 The magnetic quality of the Fe0.5Rh0.5 is only one part of the challenge. For the 9 magnetoelectric function to be produced we must be able to apply electric fields to the piezoelectric 10 PMN-PT and produce strains that drive changes in the magnetization. To gauge how the Fe0.5Rh0.5 11 deposition process (high temperature, low oxygen pressure) impacts the PMN-PT we probed the 12 room-temperature leakage behavior of several heterostructure types (Experimental Section and 13 Suppl. Fig. S5). For the Fe0.5Rh0.5/PMN-PT heterostructures, large and (fairly) symmetric leakage 14 current-voltage behavior is observed. This is consistent with the diminished magnetization data for 15 these same heterostructures and points to the fact that the high-temperature and reducing growth 16 environment for the Fe0.5Rh0.5 likely results in the oxidation of the Fe0.5Rh0.5 and deterioration of 17 the PMN-PT that ultimately limits the utility of these heterostructures. This is made more obvious 18 by comparing a test structure of room-temperature deposited platinum as the top contacts to the 19 PMN-PT which show considerably lower leakage current densities, especially under negative bias 20 (applying voltage from the top metal layer to the bottom Ba0.5Sr0.5RuO3). Such structures also 21 show highly asymmetric leakage profiles which is due to the asymmetry of the top metal (platinum) 7 1 and bottom Ba0.5Sr0.5RuO3 electrodes. To address this challenge, the MgO-buffer layer was 2 explored and leakage studies of the Fe0.5Rh0.5/MgO/PMN-PT heterostructures reveal responses 3 more akin to those of the Pt/PMN-Pt structures; in other words, the thin insulating MgO likely 4 suppresses the oxidation of the Fe0.5Rh0.5 and the deterioration of the PMN-PT during the 5 subsequent high-temperature Fe0.5Rh0.5 growth. In all, these studies of crystal structure and 6 structural quality, magnetic properties of the Fe0.5Rh0.5, and leakage behavior of the devices suggest 7 that it is possible to produce high-quality, epitaxial Fe0.5Rh0.5 thin films on PMN-PT thin films and 8 that the inclusion of a thin MgO layer both improves the Fe0.5Rh0.5 quality and suppresses 9 deterioration of the PMN-PT films during high-temperature, low-pressure Fe0.5Rh0.5 growth. 10 Having established the quality of the Fe0.5Rh0.5/MgO and Fe0.5Rh0.5/MgO/PMN-PT 11 heterostructures, we focus on these two heterostructures types as we explore methodologies to 12 probe magnetoelectric coupling in micron-scale devices. Since the Fe0.5Rh0.5 simultaneously 13 serves as the active magnetic layer and the top electrode it must be defined into small devices sizes, 14 thus making use of traditional magnetic probes (such as magnetometry) challenging. Alternatively, 15 we can leverage magnetotransport approaches and especially the anomalous Hall effect (AHE) to 16 probe the magnetoelectric coupling instead of relying on direct measurements of magnetization 17 under applied electric fields. To probe the AHE, the Fe0.5Rh0.5/MgO and Fe0.5Rh0.5/MgO/PMN-PT 18 heterostructures were fabricated into Hall bars (~5 μm x 15 μm in size; Experimental Section and 19 Fig. 3a). Before quantifying the magnetoelectric coupling, we first measured temperature- 20 dependent AHE curves for the Fe0.5Rh0.5/MgO (Fig. 3b) and Fe0.5Rh0.5/MgO/PMN-PT (Fig. 3c) 21 heterostructures. The Fe0.5Rh0.5/MgO heterostructures serve as a reference for the AHE changes 8 1 during the AFM-to-FM phase transition in Fe0.5Rh0.5. At 250 K, the Fe0.5Rh0.5 in both 2 heterostructures is in the pure AFM phase with negligible magnetization and, hence, the ordinary 3 Hall effect dominates the transport. As the temperature increases, the Hall resistivity (xy) of the 4 Fe0.5Rh0.5 becomes dominated by contributions from the AHE, which indicates that the AFM-to- 5 FM phase transition has been initiated upon increasing temperature. As a result, xy increases with 6 magnetic field and falls back with decreasing magnetic field. Both the Fe0.5Rh0.5/MgO and 7 Fe0.5Rh0.5/MgO/PMN-PT heterostructures behave similarly in nature, but with the temperature for 8 the onset of and completion of the AFM-to-FM transition being higher in the 9 Fe0.5Rh0.5/MgO/PMN-PT heterostructure. 10 As it pertains to the magnetoelectric coupling, the ultimate goal is to produce a quantitative ΔM 11 estimate of the magnetoelectric coupling coefficient αME = μ0 where M is the magnetization E 12 and E is the applied electric field. While E can be readily understood from the applied voltage and 13 device geometry, the estimation of ΔM poses a challenge due to the lack of direct measurements 14 of magnetization in the small volume/area devices. To overcome this challenge, we leveraged the 15 linear temperature-dependence of the spin-flop field (SFF, the magnetic field at which a spin-flop 16 transition occurs) extracted from the temperature-dependent Hall-resistivity studies (Fig. 3d). 17 Again, the nature of the two heterostructures is the same with nearly identical slopes for the 18 response with varying temperature. Therefore, we can use this information to estimate the ΔM 19 value under a given E by first finding the temperature which is expected to have the same SFF as 20 the SFF under the applied E. In turn, the difference in temperature is then converted to an estimated 21 ΔM using the magnetization-temperature curves from the same heterostructures (Fig. 2). 9 1 To quantify αME , negative voltages were applied to the Fe0.5Rh0.5/MgO/PMN-PT 2 heterostructures and the resulting magnetotransport was measured at 325 K (Experimental Section, 3 Fig. 4a). Of course, there is no effect in the Fe0.5Rh0.5/MgO heterostructures and similar 4 experiments were tried for the Fe0.5Rh0.5/PMN-PT, but the large leakage effects limited their utility 5 (Suppl. Fig. S6). The xy curves reveal signs of voltage dependence in that the SFF increases with 6 increasing magnitude of applied voltage (Fig. 4a), which indicates that the AFM-to-FM phase 7 transition occurs at different fields under applied voltage. Observable changes in AHE can even 8 be seen at -0.5 V (-50 kV/cm), which saturate at -0.75 V (-75 kV/cm). Using the aforementioned 9 linear relationship between SFF and temperature (Fig. 3d), the SFF changes at 325 K for the 10 Fe0.5Rh0.5/MgO/PMN-PT heterostructures under different voltages are used to estimate the 11 temperature at which the Fe0.5Rh0.5 has the same SFF. The difference between the magnetization 12 at this temperature and the magnetization at 325 K (δM) is then extracted from the magnetization- 13 temperature curves (Fig. 2e); a summary of the pertinent values of important terms (i.e., applied 14 voltage, SFF at 325 K, the equivalent temperature for the same SFF at 0 V, the estimated δM, and 15 the extracted αME ) are provided (Fig. 4b). The extracted αME as a function of negative voltage 16 shows a sharp increase between -0.5 and -0.75 V and reaches a value of 5.29 x 10-9 s/m at 325 K 17 at the maximum (Fig. 4c). At higher voltages, further changes in the AHE are not observed, likely 18 due to the potential leakage at high electric fields and/or saturation of the piezoelectric strain of 19 the films. 20 The observed values of αME can be compared to prior works on bulk crystals and thin films 21 (Fig. 5). Ideally, the combination of low voltage (< 100 mV) actuation, room-temperature (300 K) 10 1 functionality, and resulting large αME characteristics are desired. As the large scatter in the data 2 suggests, however, this is a challenging task. Although Fe0.5Rh0.5/MgO/PMN-PT heterostructures 3 satisfy the low voltage (-0.75 V) and near room-temperature (325 K) characteristics, they suffer 4 from relatively small αME , which are likely ultimately due to the reduced piezoelectric strains that 5 are possible in clamped films36 and due to the thin (relatively lower dielectric constant) MgO- 6 buffer layer which is required to address issues in integrating the metal ferromagnetic layer with 7 PMN-PT across which some of the voltage is dropping. To demonstrate the potential of further 8 improvements in thin films, we estimate the effective αME based on the actual voltage drop across 9 PMN-PT layer. Here, the actual voltage drops in MgO (dielectric permittivity of 8) and PMN-PT 10 (dielectric permittivity of 1000) layer can be estimated from charge continuity conditions at the 11 interface.34 Using these values, when -0.75 V is applied to the device, the voltage drops across the 12 MgO and PMN-PT layers are estimated to be -0.694 V and -0.056 V, respectively. Using this rather 13 simple estimation, one obtains an effective αME of 7.08 x 10-8 s/m at 325 K, which are comparable 14 to αME obtained from using single-crystal BaTiO3 and 0.72PMN-0.28PT substrates.37,38 We expect 15 that progress in the fabrication of laterally defined structures (thereby removing some of the in- 16 plane clamping of the films) and the integration of piezoelectric membranes released from the 17 substrate39,40 and improvements in the growth process41 can lead to further refinements of the 18 magnetoelectric response. 19 20 CONCLUSIONS 21 In summary, it has been shown that it is possible to translate similar magnetoelectric 11 1 heterostructures studied in the bulk to the thin film regime. It is possible to produce high-quality, 2 epitaxial, (001)-oriented films of Fe0.5Rh0.5 on PMN-PT and MgO/PMN-PT thin films. Subsequent 3 temperature-dependent magnetization and resistivity studies reveal the ability to produce near- 4 room-temperature phase transitions in the Fe0.5Rh0.5. A methodology to probe and quantify the 5 magnetoelectric coupling in small thin-film devices via anomalous Hall effect studies was also 6 developed wherein one can compare the spin-flop field change induced by temperature and 7 external voltage to extract a measure of the magnetoelectric coupling coefficient. Using this 8 approach, αME was estimated to be 5.29 x 10-9 s/m at 325 K under -0.75 V and as large as 7.08 x 9 10-8 s/m if one accounts for the voltage drop only across the PMN-PT. While this is promising for 10 next-generation electric-field control of magnetization, improvements in how we integrate the 11 dissimilar materials are needed and nanofabrication approaches to enable large piezoelectric 12 responses in thin films are needed. 12 1 EXPERIMENTAL SECTION 2 Epitaxial thin-film synthesis. Pulsed-laser deposition (PLD) using a KrF excimer laser 3 (248 nm, LPX 300, Coherent) was used to grow 100 nm 0.68Pb(Mg1/3Nb2/3)O3-0.32PbTiO3 (PMN- 4 PT) / 25 nm Ba0.5Sr0.5RuO3 (BSRO) heterostructures on NdScO3 (NSO) (110) substrates (CrysTec 5 GmbH). The PMN-PT growth was carried out at a heater temperature of 600 C in a dynamic 6 oxygen pressure of 200 mTorr with a laser fluence of 1.8 J/cm2 and a laser repetition rate of 2 Hz 7 from a ceramic target (Praxair) of the same composition with 10% lead excess to compensate for 8 lead loss during growth. The BSRO growth was carried out at temperature of 750 C in a dynamic 9 oxygen pressure of 20 mTorr with a laser fluence of 1.85 J/cm2 and a laser repetition rate of 3 Hz 10 from a ceramic target (Praxair) of the same composition. Following the growth, the samples were 11 cooled to room temperature at 5C/min. in a static oxygen pressure of 700 Torr. In the case of 12 MgO-buffered heterostructures, 3-10 nm MgO buffer layer was grown by PLD immediately after 13 PMN-PT at the PMN-PT growth temperature of 600 C in a dynamic oxygen pressure of 20 mTorr 14 with a laser fluence of 2.5 J/cm2 and a laser repetition rate of 15 Hz from a ceramic target (Praxair) 15 of the same composition. Fe0.5Rh0.5 films were grown using DC sputtering techniques. For the 16 sputtering deposition of Fe0.5Rh0.5, the PMN-PT / BSRO / NSO heterostructures were heated to 17 500-700 °C in a base pressure of at least 10-8 Torr. Subsequently, a small amount of argon gas was 18 introduced into the chamber, maintaining the total pressure at 3.0 mTorr. The Fe0.5Rh0.5 films were 19 then grown from a stoichiometric Fe0.5Rh0.5 target. After the deposition, films were annealed at 20 various temperatures between 600 and 800 °C in vacuum for 40 min. and subsequently cooled to 21 room temperature at 10 °C/min. 3 nm thick capping platinum layers were also grown by DC 13 1 sputtering at room-temperature and a target power of 100 W in a dynamic argone pressure of 2 2 mTorr to prevent the oxidation of FeRh. 3 Structural characterization. X-ray θ-2θ line scans about symmetrical reflections of the 4 films and substrates were conducted with a high-resolution X-ray diffractometer (X'pert Pro2, 5 PANalytical). 6 Magnetometry. Magnetic characterizations were carried out with a superconducting 7 quantum interference device magnetometer (Quantum Design, 2 K, 7 T), with the magnetic field 8 applied along the out-of-plane direction of the film. 9 Transport measurements. The temperature-dependent resistance and Hall effect 10 measurement were carried out in a CRYOGENIC measurement system (2 K, 14 T) with the 11 magnetic field applied perpendicular to the device plane. A constant 10 𝜇𝐴 dc current was applied 12 for all temperature-dependent resistivity measurements. A constant 100 𝜇𝐴 dc current was applied 13 for all Hall measurements with a series of back gating voltages. 14 Device fabrication for AHE studies. To measure the physical properties of the Fe0.5Rh0.5 / 15 (MgO) / PMN-PT heterostructures, the active region was patterned into a 6-point contact Hall bar 16 geometry (20 µm × 400 µm) by the following fabrication process. First, photoresist was patterned 17 by spin-coating on the blanket Pt/ Fe0.5Rh0.5 / (MgO) / PMN-PT heterostructures. Second, the 18 photoresist was patterned into Hall-bar geometries via standard optical photolithography. Third, 19 the heterostructures were ion-milled to define the active region. Fourth, a 300-nm-thick insulating 20 layer of MgO was deposited at room-temperature by PLD to separate the bottom contacts from the 21 top contacts. Fifth, the photoresist was removed in acetone and a second photoresist layer was 22 spin-coated for the second photolithography step. Sixth, inverse patterns of platinum contact pads 23 were formed by photolithography. Seventh, a 100-nm-thick layer of platinum was deposited to 14 1 form contact pads. Eighth, the photoresist was removed in acetone. Ninth, the heterostructures 2 were mounted on a chip carrier and wire-bonded for insertion in the measurement system. 3 4 ACKNOWLEDGEMENTS: 5 W.Z. and J.K. contributed equally to the work. W.Z., J.K., and L.W.M. acknowledge the Intel Corp. 6 via the FEINMAN program. X.H and R.R. acknowledge the support of ASCENT, one of the SRC- 7 JUMP Centers. D.P. acknowledges funding from the European Union's Horizon 2020 research and 8 innovation program under the Marie Skłodowska-Curie grant agreement No. 79712. G.A.P.V. 9 acknowledges support from the National Science Foundation under grant DMR-1708615 and from 10 the NSF GRFP 15 1 Figures/Captions 2 3 Figure 1. Schematic illustrations of the three heterostructure variants used and described in this 4 work including the (a) Fe0.5Rh0.5/MgO, (b) Fe0.5Rh0.5/PMN-PT, and (c) Fe0.5Rh0.5/MgO/PMN-PT 5 structures. (d) θ-2θ X-ray diffraction scans for the same three heterostructure variants revealing 6 the ability to produce fully epitaxial, 00l-oriented films of Fe0.5Rh0.5 in all cases. (e) Rutherford 7 backscattering spectrometry spectra for an as-grown Fe0.5Rh0.5/MgO heterostructure at the 8 optimized growth condition. 9 16 1 Figure 2. Magnetization and resistivity data as a function of temperature for the as-grown 2 heterostructures, including (a) magnetization and (b) transport data for the Fe0.5Rh0.5/MgO 3 heterostructures, (c) magnetization and (d) transport data for the Fe0.5Rh0.5/PMN-PT 4 heterostructures, and (e) magnetization and (f) transport data for the Fe0.5Rh0.5/MgO/PMN-PT 5 heterostructures. 6 17 1 Figure 3. (a) Schematic illustration of the Hall-bar devices fabricated into the various 2 heterostructures. Anomalous Hall resistivity as a function of applied magnetic field at various 3 temperatures for the (b) Fe0.5Rh0.5/MgO and (c) Fe0.5Rh0.5/MgO/PMN-PT heterostructures. (d) 4 Summary of the spin-flop field (SFF) as a function of temperature for the different heterostructure 5 variants revealing similar temperature-dependence for both heterostructures. 6 18 1 Figure 4. (a) Anomalous Hall resistivity as a function of magnetic field measured at different 2 applied dc voltages for the Fe0.5Rh0.5/MgO/PMN-PT heterostructures reveals a strong 3 magnetoelectric effect with a sharp raise in magnitude between -0.5 and -0.75 V. (b) Summary of 4 important values from both the magnetization and transport studies to estimate the magnetoelectric 5 coupling. (c) Evolution of the magnetoelectric coupling coefficient αME as a function of applied 6 voltage in this system. 7 19 1 Figure 5. (a) Graphical and (b) table summary of studies of magnetoelectric coupling coefficient 2 αME in ferromagnetic/piezoelectric-ferroelectric heterostructures (including both bulk and thin- 3 film forms). The ultimate goal is a combination of low-voltage, room-temperature, and high αME 4 which remains challenging to obtain. Pertinent references include: 1) R. Cherifi, et al. Nature 5 Mater. 13, 345–351 (2014); 2) J. Heron, et al. Nature 516, 370–373 (2014).; 3) C. Thiele, et al. 6 Phys. 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