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Tatsuya Yamamoto, Tomohiro Ichinose, [Jun Uzuhashi](https://orcid.org/0000-0003-2023-8158), Takayuki Nozaki, [Tadakatsu Ohkubo](https://orcid.org/0000-0003-3548-1951), Kay Yakushiji, Shingo Tamaru, Hitoshi Kubota, Shinji Yuasa

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[Magneto-transport properties in perpendicularly magnetized magnetic tunnel junctions using an Mg40Fe10O50 tunnel barrier](https://mdr.nims.go.jp/datasets/bf1aa179-72b3-4dca-90e3-8bc155c6fa0c)

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Magneto-transport properties in perpendicularlymagnetized magnetic tunnel junctions using anMg40Fe10O50 tunnel barrierTatsuya Yamamoto∗, Tomohiro Ichinose, Jun Uzuhashi, Takayuki Nozaki,Tadakatsu Ohkubo, Kay Yakushiji, Shingo Tamaru, Hitoshi Kubota, ShinjiYuasaNational Institute of Advanced Industrial Science and Technology (AIST), Research Centerfor Emerging Computing Technologies, Tsukuba, Ibaraki 305-8568, JapanNational Institute for Materials Science (NIMS), Tsukuba, Ibaraki 305-0047, JapanAbstractWe develop perpendicularly magnetized magnetic tunnel junctions (MTJs)consisting of a CoFeB/Mg40Fe10O50 (MgFeO)/CoFeB multilayer. The use ofMgFeO yields a substantial improvement in the flatness of the MTJ film stackcompared with conventional MTJ films using MgO as a tunnel barrier layer,and a 1.7 times enhancement in the perpendicular magnetic anisotropy energyis obtained for the ultrathin CoFeB layer deposited on the MgFeO layer. Nanos-tructural analysis combined with elemental distribution mapping reveals theformation of highly (001)-oriented MgFeO in the as-deposited MTJ film, andthe crystalline MgFeO layer effectively inhibits diffusion of B atoms from theCoFeB layers through the tunnel barrier layer during the post-annealing pro-cess. Accordingly, the MgFeO-MTJ exhibits exceptional stability against hightemperature annealing, and a large tunnel magnetoresistance ratio of 235%is demonstrated in MTJ nanopillar devices after annealing at 400◦C. Ferro-magnetic resonance measurements also reveal a reduced magnetic damping inthe MgFeO-MTJs owing to the improved uniformity in the CoFeB layer. Thepresent experimental results will facilitate the development of magnetoresistive∗Corresponding authorEmail address: yamamoto-t@aist.go.jp (Tatsuya Yamamoto)Preprint submitted to Acta Materialia January 19, 2024memory devices with increasing memory density and higher energy efficiencies.Keywords: perpendicular magnetic anisotropy, magnetic tunnel junction,magnetoresistive random access memory1. IntroductionMagnetic tunnel junctions (MTJs) consisting of a ferromagnetic layer/tunnelbarrier/ferromagnetic layer junction are fundamental elements in the field ofspintronics. The discoveries of large tunnel magnetoresistance (TMR) [1, 2, 3, 4,5, 6, 7, 8] and large perpendicular magnetic anisotropy (PMA) in CoFe(B)/MgO/CoFe(B)5junctions [9, 6, 8, 10, 11, 12] have led to the rapid development of magnetoresis-tive random access memory (MRAM) devices [13, 7, 14, 15, 16]. In the MRAMcell, binary data are stored in one of the ferromagnetic layers (called the freelayer) in the MTJ by electrically controlling the magnetization direction, whilethe other ferromagnetic layer possesses a fixed magnetization to serve as a ”ref-10erence” for the magnetization direction of the free layer [17]. To increase thememory density while reducing the energy required to switch the magnetiza-tion, a large PMA as well as a low magnetic damping are required for thefree layer [13, 7, 14]. The MRAM application also requires the MTJ devicesto be formable on Si substrates with a polycrystalline electrode using a mass-15production-compatible thin film deposition process (e.g., magnetron sputtering)and to be tolerant of high temperature (> 400◦C) annealing to overcome theCMOS back-end-of-line process [15].Several core stacking structures to fulfill the above requirements, such as asynthetic antiferromagnetic layer based on CoPt artificial superlattices [18], an Ir20spacer layer exhibiting strong interlayer exchange coupling [19], and MgO/CoFeB/X/CoFeB/MgOmultilayered free layer [20, 21, 22]. In particular, the stacking structure ofMgO/CoFeB/X/CoFeB/MgO has been intensively studied towards improvingthe TMR and PMA after high temperature annealing [23, 21, 22, 24, 25, 26].However, the use of an MgO/CoFeB/X/CoFeB/MgO free layer imposes another25difficulty in preparing flat and uniformly magnetized CoFeB layers; even though2the CoFeB thin films favor an amorphous phase without crystal grains, the poorwettability of CoFeB on MgO leads to island growth of the CoFeB layer [27].Such structural inhomogeneity in the free layer is unfavorable in terms of energyefficiency for the electrical control of magnetization since the structural inhomo-30geneity inevitably leads to a magnetic inhomogeneity in the free layer and thusa larger energy dissipation during the magnetization switching. The impact ofstructural inhomogeneity can be even more severe for MRAM devices driven bythe voltage-controlled magnetic anisotropy (VCMA) effect at the MgO/CoFe(B)interface [28, 29, 30, 31, 32], in which the free layer is typically as thin as 1 nm35to minimize the electric-field screening effect. A possible solution to reducingthe structural inhomogeneity is to lower the substrate temperature during thedeposition of the CoFeB layer. In our previous studies, we have shown that thedeposition of CoFeB layers at a cryogenic substrate temperature enhances thePMA as well as TMR in the MgO/CoFeB junction while reducing the extrinsic40component of magnetic damping due to magnetic inhomogeneity [33, 34]. An-other solution may be to insert an ultrathin γ-Fe2O3 layer in between the MgOand CoFeB layers, as the ultrathin γ-Fe2O3 insertion layer has been demon-strated to improve the wettability of Fe on MgO without affecting the TMRratio [35].45In this work, we develop perpendicularly magnetized MTJs using an Mg40Fe10O50(MgFeO) tunnel barrier and a CoFeB layer deposited at a cryogenic tempera-ture. In contrast to Ref. [35], in which an MgO/γ-Fe2O3 bilayer is used astunnel barrier, here we utilize the interfacial segregation of Fe atoms from apartially Fe-substituted MgO layer [36, 37] to improve the wettability of CoFeB.50The use of MgFeO barrier layer substantially improves the homogeneity of freelayer, which contributes to an enhanced PMA and reduces effective magneticdamping. More interestingly, the MgFeO-MTJs exhibit higher tolerance againsthigh-temperature annealing, and a large TMR ratio of 235% is achieved afterannealing at 400◦C for 2h. We investigate the improved magnetic and electrical55transport properties in MgFeO-MTJs through detailed magnetic measurementsand nanostructural analysis for the MTJ films as well as magneto-transport3measurements on MTJ nanopillar devices.2. ExperimentalThin films consisting of CoFeB (0.8 nm)/MgO or MgFeO (2.5 nm)/CoFeB60(0.7- 1.0 nm)/Mo (0.3 nm)/CoFeB (0.6 nm)/MgO (1.0 nm) MTJ stacks weredeposited on 300 mm Si wafers with a buffer layer using an ultrahigh-vacuummagnetron sputtering system (EXIM) manufactured by Tokyo Electron Ltd.The CoFeB/Mo/CoFeB multilayer acts as a free layer whose magnetization di-rection can be switched by a relatively small external field, whereas the bottom65CoFeB layer magnetization is pinned by a CoPt/Ir-based synthetic antiferro-magnet to serve as a reference layer. The MgO layer was deposited by DCsputtering of Mg followed by in-situ post oxidation using a pure O2 gas whereasthe MgFeO layer was deposited by RF sputtering from an Mg40Fe10O50 tar-get. We also fabricated MgO-MTJs in which the MgO layer is deposited by RF70sputtering from an MgO target and confirmed that the fabrication process ofthe MgO layer does not affect the magnetic and electric transport properties ofthe MgO-MTJs. The resistance-area product (RA) of MTJs used in this workis around 500 ohm·µm2, which is compatible with the VCMA studies. For bothMgO- and MgFeO-MTJs, the CoFeB layer on top of the tunnel barrier layer75was deposited at −173◦C, followed by the deposition of Mo/CoFeB bilayer atroom temperature. The magnetic properties of unpatterned films were charac-terized by using a vibrating sample magnetometer (VSM) and a ferromagneticresonance measurement system equipped with a network analyzer (VNA-FMR)[38]. The nanostructure and elemental distribution of the MTJ stacks were in-80vestigated by scanning transmission electron microscopy (STEM), nano-beamelectron diffraction (NBD), electron energy loss spectroscopy (EELS), and en-ergy dispersive X-ray spectroscopy (EDS) using FEI Titan G2 80-200. Thinfoil specimens for the nanostructural analyses were prepared by a standardlift-out method using a focused ion beam with a scanning electron microscopy85system (FEI Helios G4UX). To characterize the magneto-transport properties,4Figure 1: Layer structures and magnetic properties of MTJ films: (a,c) ADF-STEM images,(b,d) EDS elemental maps, (e) Major magnetization curves, (f) Minor magnetization curves ofMgFeO-MTJ with different ttotal, and (g) µ0Mst as a function of ttotal. The MTJ films wereannealed at Ta = 400◦C for 1h after deposition. Lines in (g) denote linear fits to experimentaldata.the MTJ films were microfabricated into nanopillar devices with a diameterof approximately 75 nm using electron-beam lithography and Ar-ion etching.The resistance-field (R-H) and differential conductance G were measured byconnecting the device to a source measurement unit (Keysight B2902A) and90a lock-in amplifier (Stanford Research Systems SR865), respectively. All mag-netic and electrical measurements presented in this article were done at roomtemperature.53. Results and discussion3.1. Layer morphology, nanostructure and atomic distribution95Figures 1(a) and (b) show annular dark field (ADF) STEM images of MTJfilms using MgO and MgFeO tunnel barriers, respectively. Both films were ex-situ annealed in a vacuum at a temperature Ta of 400◦C for 1h. As expectedfrom the poor wettability of CoFeB on MgO, the contrast at the MgO/CoFeBinterface is blurred for the MgO-MTJ. Meanwhile, the ADF-TEM image taken100of the MgFeO-MTJ reveals the formation of a flat MgFeO/CoFeB interface.The EDS elemental mapping shown in Figs. 1(c) and (d) also reveal improvedflatness of the CoFeB free layer in the MgFeO-MTJ. Figure 1(e) shows the mag-netization curves of MTJ films obtained under an out-of-plane H. After theannealing at Ta = 400◦C, both MTJs exhibit perpendicular magnetization. In-105terestingly, we found that the use of MgFeO barrier not only affects the PMA inthe free layer but also alters the magnetic stability of the reference layer coupledwith the CoPt/Ir-based SAF layer. The possible origin of improved magneticproperties of reference layer and CoPt/Ir-based SAF layer will be discussed indetail in a later section. Figure 1(f) shows the minor magnetization curves of the110MgFeO-MTJ films with different total CoFeB layer thickness ttotal. Here, ttotalcorresponds to the thickness of the CoFeB/Mo/CoFeB free layer excluding theMo layer thickness. For this thickness range, the CoFeB/Mo/CoFeB free layerexhibits distinct PMA and the magnetization direction can be sharply switchedby applying a small magnetic field. The coercivity slightly increases with de-115creasing ttotal, which can be attributed to an increase in the effective PMAenergy density, Keff , as will be revealed later by the FMR measurement. Fig-ure 1(g) shows the saturation magnetization (µ0Mst) of the CoFeB/Mo/CoFeBfree layer obtained from the minor magnetization curves as a function of ttotal.For ttotal ≥ 1.4 nm, the µ0Mst value almost linearly increases with increasing120ttotal for both the MgFeO- and MgO-MTJs, and the linear fit to the experi-mental data gives a dead layer thickness, tdead = 0.34 nm (0.39 nm), as well asµ0Ms = 2.03 (T) (1.97 T) for the MgFeO-MTJ (MgO-MTJ). The formation of6Figure 2: Nanostructural analysis on the MTJ films: (a, b) High-resolution BF-STEM imagesand NBD patterns for each layer, and (c, d) Virtual ADF-STEM images using the selected(002) and (002̄) diffraction spots as indicated by the dotted circles (a) and (b).dead layer can be explained by the intermixing of CoFeB and Mo as well as theoxidation of CoFeB surface during the deposition of MgO capping layer. The125slight reduction in tdead in the MgFeO-MTJ is attributed to the segregation offerromagnetic Fe from the MgFeO barrier associated with annealing [36].Figures 2(a) and (b) show high-resolution bright field (BF) STEM imagesalong with NBD patterns obtained from the MgO- and MgFeO-MTJs annealedat 400◦C, respectively. In accordance with previous studies [39, 3, 40], the130BF-STEM images as well as the NBD patterns reveal the formation of (001)-oriented MgO and MgFeO layers, and the CoFeB layers crystallize from theMg(Fe)O/CoFeB interface. To compare the crystallinity of the Mg(Fe)O layer,virtual ADF-STEM images [41] are reconstructed using only the selected (002)and (002̄) diffraction spots as indicated by the dotted circles in Figs. 2(a) and135(b) and shown in Figs. 2(c) and (d). The contrast of these images represents thediffraction intensity at the selected regions in the NBD patterns, i.e., the brighterareas contain a larger crystalline component having the same crystalline plane.In the MgO-MTJ, although a definite (001) orientation of MgO layer is observed7Figure 3: Cross-sectional images of MTJ films and compositional analyses: (a-c) BF-STEMimages (d-e) EELS maps showing the distribution of B atoms. (g-f) Compositional line profilesobtained from the EDS and EELS analyses.in both the BF-STEM image and the NBD patterns, the crystalline plane is140found to be non-uniform over the film plane and thus contains many grainboundaries. In contrast, the reconstructed MgFeO layer is almost continuousalong the film plane, which indicates good in-plane crystalline uniformity of theMgFeO tunnel barrier layer.Figures 3(a)-(c) show BF-STEM images of as-deposited MgFeO-MTJ, MgFeO-145MTJ annealed at 400◦C, and MgO-MTJ annealed at 400◦C, respectively, andFigs. 3(d)-(f) illustrate the distribution of B atoms in the corresponding areaobtained from the EELS analysis. As shown in Fig. 3(a), the MgFeO layer8exhibits good crystallinity even in the as-deposited film, which can effectivelyeliminate the atomic diffusion from the CoFeB layer. In fact, the EELS maps150shown in Figs. 3(d) and (e) reveal that the B atoms in the as-deposited CoFeBlayers are confined in the CoFeB layers even after the annealing at 400◦C. In theMgO-MTJ, however, the B atoms are dispersed from the CoFeB layer into theSAF and the capping layers, which suggests the diffusion of B atoms throughthe grain boundaries in the MgO layers. The integrated EELS intensities are155shown in Figs. 3(g)-(i) together with compositional line profiles obtained fromthe EDS analysis. In the MgFeO-MTJ, the B atoms move away slightly fromthe MgFeO/CoFeB interface after annealing at 400◦C associated with the crys-tallization of CoFeB to CoFe, but most B atoms remain in the CoFeB layers. Inthe MgO-MTJ annealed at 400◦C, the EELS intensity at the CoFeB/Mo/CoFeB160free layer is relatively weak and there are small peaks in the CoPt layer as wellas the capping layer as denoted by the dotted circles in Fig. 3(i). Thus, thedegraded magnetic property of SAF in the MgO-MTJ as observed in Fig. 1(e)can be explained by the diffusion of B atoms to the CoPt layer associatedwith annealing. The EDS line profiles also reveal the segregation of Fe from the165MgFeO layer after the annealing; in the as-deposited MgFeO-MTJ, about 9 at%Fe is found in the MgFeO layer, whereas the Fe content is reduced to below 5at% in the annealed MgFeO-MTJ and the diffused Fe atoms concentrate at theMgFeO/CoFeB interfaces. As a result, the MgFeO/CoFeB interfaces becomeFe-rich compared with MgO/CoFeB interfaces in the MgO-MTJ. It is notewor-170thy that the atomic distributions in the CoFeB/Mo/CoFeB/MgO multilayer onthe MgFeO layer are narrow and the transitions at the interfaces are sharperthan those in the MgO-MTJ, which verifies the improved free layer flatness inthe MgFeO-MTJ.3.2. Magnetic anisotropy and magnetic damping of free layer175Figure 4(a) shows FMR spectra obtained from the MTJ films under an out-of-plane H. These spectra were obtained by sweeping H under a constant rfexcitation at a frequency f of 35 GHz. The total CoFeB thickness ttotal is 1.59Figure 4: Results of FMR measurements: (a) FMR spectra obtained from the MTJ filmsunder an out-of-plane H with f = 35 GHz, (b) Relationship between Hr and f. (c) Keff teffas a function of teff . Lines in (b) and (c) denote linear fits to the experimental data.nm and the films were annealed at 400◦C. For the same ttotal, the free layerof the MgFeO-MTJ exhibits a narrower resonance linewidth ∆H and a larger180resonance signal than that of the MgO-MTJ, suggesting a larger magnetizationand a smaller magnetic damping in the MgFeO-MTJ. We also found that theresonance field (Hr) of the free layer in the MgFeO-MTJ is higher than that ofMgO-MTJ under the same f . The relationship between Hr and f is shown inFig. 4(b). For both MTJs, there is a linear relationship between Hr and f , and185the anisotropy field Hk of the free layers are obtained from linear fits to theexperimental data:f =gµ0µBh(Hr −Hk), (1)10where h is Planck’s constant, g is the Lande g-factor, and µB is the Bohr mag-neton. Then, the areal PMA energy density Keffteff (teff = ttotal − tdead) isestimated by using the obtained Hk values: Keffteff = MsHk/2. Figure 4(c)190shows the obtained Keffteff values as a function of teff for both MTJ films. TheMgFeO-MTJs exhibit larger Keffteff than MgO-MTJs regardless of teff , and amaximum Keffteff of 0.69 (0.41) mJ/m2 is obtained at teff = 0.96 (1.0) nm forthe MgFeO-MTJs (MgO-MTJs). In addition, the y-intercept of the dependenceof Keffteff on teff yields an estimated interfacial magnetic anisotropy energy195density, Ki, of 2.1 and 1.5 mJ/m2 for the MgFeO- and MgO-MTJs, respec-tively. The enhanced interfacial PMA in the MgFeO-MTJ is attributed to theimproved free layer flatness as well as the formation of Fe-rich region at theMgFeO/CoFeB interface [27].Figures 5(a) and (b) show the FMR linewidth ∆H of the free layer as a200function f for the MgFeO- and MgO-MTJs, respectively. For the both films,∆H increases as ttotal is reduced. To investigate the origin of increased ∆H indetail, we applied linear fitting to the experimental data [42]:∆H =2hαtotalgµ0µBf +∆H0, (2)where ∆H0 is the inhomogeneous linewidth originating from the structuraland/or magnetic inhomogeneities in the free layer. αtotal corresponds to the205total magnetic damping which includes the radiative damping [43] and the in-terfacial enhancement of damping due to the spin pumping effect [44, 45, 42].However, we neglect the radiative damping in the following discussion since itscontribution can be ignored for ultrathin films. Note that we also omit the dataobtained from the MgO-MTJ with ttotal = 1.3 nm, because there is an anoma-210lous enhancement of ∆H for ttotal = 1.3 nm as shown in the inset of Fig. 5(b),which may originate from discontinuity in the free layer.The ∆H0 and αtotal values obtained from the fitting to the experimental datausing Eq. (1) are shown in Fig. 5(c) and (d), respectively. For the same teff ,the MgFeO-MTJ exhibits a smaller ∆H0 compared with the MgO-MTJ. The215reduced ∆H0 in the MgFeO-MTJ can be attributed to the improved flatness11Figure 5: FMR linewidth and magnetic damping of free layers with various ttotal: (a, b) ∆Has a function of f , (c) ∆H0 as a function of teff , and (d) αtotal as a function of 1/teff . Linesin (d) denote linear fits to the experimental data.of the free layer as determined by the TEM-EDS analyses. As is the case with∆H0, αtotal increases with decreasing teff . The dependence of αtotal on 1/teffis shown in Fig. 5(d). The y-intercept corresponds to the intrinsic magneticdamping, αint, which is free of the spin pumping effect and the value should be220the same for ferromagnets with the same composition and crystal structure. Theobtained αint values, αint = 0.0055 for the MgFeO-MTJ and αint = 0.0072 forthe MgO-MTJ, are in agreement with those reported for FeCo alloy thin films[42], while the shallower slope of αtotal on 1/teff for the MgFeO-MTJ suggestsa reduced magnon scattering owing to the improved film uniformity and/or22512Figure 6: Magneto-transport properties of MTJ nanopillar devices: (a) R-H curves measuredunder an out-of-plane H with V = 10 mV. (b-d) V dependence of (b) TMR(V)/TMR(10mV), (c) GAP(V )/GAP(0), and (d) GP(V )/GP(0). ttotal = 1.5 nm for both devices.reduced atomic interdiffusion.3.3. Magneto-transport properties of MTJ nanopillarsFinally, we discuss the magneto-transport properties of nanopillar MTJ de-vices. Figure 6(a) shows R-H curves of MgFeO- (filled circles) and MgO-MTJdevices (open circles) measured under a perpendicular H and a dc bias voltage230V of 10 mV. The total free layer thickness ttotal = 1.5 nm. The MTJ films wereannealed at Ta = 400◦C for 1h prior to the microfabrication. Both MTJ devicesexhibit sharp changes in resistance at around µ0H = ±0.2 T associated withthe switching of free layer magnetization, where the high- and low-resistancestate corresponds to the antiparallel (AP) and parallel (P) alignment of the free235layer and the reference layer magnetizations, respectively. The two MTJ de-vices exhibit similar AP and P resistances, and TMR ratios of around 200% areobtained for both devices. In accordance with the VSM and VNA-FMR mea-surement results, the MgFeO-MTJ device exhibits a higher magnetic stabilitythan the MgO-MTJ devices, i.e., the free layer exhibits a larger coercivity and24013Figure 7: TMR curves of MTJ nanopillar devices microfabricated from MTJ films annealedat different conditions: (a) Ta = 400◦C for 2h and (b) Ta = 420◦C for 1h. ttotal = 1.5 nm forall devices.the reference layer maintains a flat and a broader plateau for the AP state.Figures 6(b)-(d) show the V dependence of the TMR ratio and differentialconductances under the AP (GAP) and P (GP) states, respectively. The valuesof the TMR ratio and differential conductances are divided by the values atV = 10 mV and 0 mV, respectively. The use of the MgFeO tunnel barrier245alters the symmetry of conductance curve for the P state, and the deviationfrom parabolic curve at the low V region becomes more visible. These resultsare indicative of the enhanced contribution of tunneling conductance specificto crystalline CoFe/MgO/CoFe MTJs with clean interfaces [46] in the MgFeO-MTJ. It is also shown that the use of the MgFeO tunnel barrier weakens the250dependence of the TMR ratio on V . The V range for which the MTJ retainsTMR(V )/TMR(10 mV) > 0.5 is enhanced from −0.80 V ≤ V ≤ 0.65 V to −0.94V ≤ V ≤ 0.79 V by replacing the MgO tunnel barrier with the MgFeO one.The overall improvements in the magnetic and electrical transport properties inthe MgFeO-MTJ would be worth exploring in terms of practical applications.255To further investigate the tolerance of MTJs against high-Ta annealing, weprepared two sets of MTJ nanopillar devices, one annealed at Ta = 400◦C for142h [Fig. 7(a)] and the other annealed at Ta = 420◦C for 1h [Fig. 7(b)] prior tothe microfabrication. The total free layer thickness ttotal = 1.5 nm. After theannealing at Ta = 400◦C for 2h, the TMR ratio of MgFeO-MTJ is enhanced260by up to 235%. Although the TMR ratio of the MgO-MTJ is not affectedby the longer annealing time, the reduction in the coercivity and the plateauregion for the AP state indicate the degradation of PMA in both the free layerand the reference layer. The degradation of PMA is greater for the MgO-MTJannealed at Ta = 420◦C for 1h; the coercivity is reduced to as low as 0.05 T,265and the AP magnetization configuration is incomplete at H = 0 T. Meanwhile,the MgFeO-MTJ exhibits coercivity as large as 0.15 T, and the plateau of theAP state extends up to ∼ 0.3 T even after the annealing at Ta = 420◦C. Theseexperimental results clearly demonstrate that the MgFeO-MTJ is tolerant of ahigher Ta for a longer duration.2704. ConclusionsWe have evaluated the magnetic and electrical transport properties of MTJsusing MgO and MgFeO tunnel barriers. The use of MgFeO barrier substan-tially improved the flatness of the free layer owing to the improved wettabil-ity of CoFeB, and Keffteff is enhanced by a factor of 1.7 compared with that275of MgO-MTJ. At the same time, the improved homogeneity in the free layercontributed to reducing µ0∆H0 and αtotal. Detailed nanostructural analysisrevealed the uniform (001)-oriented growth of MgFeO barrier layer, which ef-fectively inhibited the segregation of B atoms from the CoFeB layers during thehigh-temperature annealing process. The improved thermal tolerance in the280MgFeO-MTJ makes it possible to attain a large TMR ratio of 235% after theannealing at Ta = 400◦C for 2h. These experimental demonstrations should bebeneficial for the development of high-density and energy-efficient MRAMs.15AcknowledgementThe authors would like to thank T. Nozaki, M. Konoto, A. Sugihara, S.285Tsunegi, Y. Hibino, L. Sakai, K. Ohba, H. Ohmori, Y. Higo, Y. Kageyamaand M. Hosomi for their fruitful discussions, and M. Toyoda, and K. Suzukifor assisting with the experiments. This work is based on results obtainedfrom a project, JPNP16007, commissioned by the New Energy and IndustrialTechnology Development Organization (NEDO), Japan.290References[1] S. S. P. Parkin, C. Kaiser, A. Panchula, P. M. Rice, B. Hughes, M. Samant,S.-H. Yang, Giant tunneling magnetoresistance at room tempereature withMgO (100) tunnel barriers, Nat. Mater. 3 (2004) 862.[2] S. Yuasa, T. Nagahama, A. Fukushima, Y. Suzuki, K. Ando, Giant room-295temperature magnetoresistance in single-crystal Fe/MgO/Fe magnetic tun-nel junctions, Nat. Mater. 3 (2004) 868.[3] D. Djayaprawira, K. Tsunekawa, M. Nagai, H. Maehara, S. Yamagata,N. Watanabe, S. Yuasa, Y. Suzuki, K. Ando, 230% room-temperature mag-netoresistance in CoFeB/MgO/CoFeB magnetic tunnel junctions, Appl.300Phys. Lett. 86 (2005) 092502.[4] J. Mathon, A. 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