Toward the applications to high-performance MRAMs [1, 2], spin-orbit torque (SOT)-induced magnetization switching has been widely investigated in various material systems. While 4d transition metals, such as Ru, exhibit low SOT due to the small spin-orbit coupling [3], recent works have reported the SOT generation in antiferromagnetic RuO2 film thanks to its unique band structure [4], although the detailed understanding is still lacking. Herein, we fabricate the (100)-oriented RuO2 film and evaluate SOT in a RuO2/Co-Fe-B bilayer by harmonic Hall measurement. 
RuO2 (4 nm)/Co20Fe60B20 (1.2 nm) bilayer was fabricated on a α-Al2O3 (0001) substrate by DC/RF sputtering, where (100)-oriented RuO2 film with a likely existence of a three-domain structure is expected to be grown [5]. The resistivity of the RuO2 film was estimated as 247.4 mΩcm from the sheet resistance. A Ru (4 nm)/Co20Fe60B20 (1.2 nm) bilayer was also fabricated for comparison. The harmonic Hall measurement was performed for Hall bar devices in a rotating magnetic field [Fig. 1]. A clear second harmonic signal is observed for RuO2/Co-Fe-B while it is negligible for Ru/Co-Fe-B, suggesting the SOT generation owing to the unique band structure of RuO2. For RuO2/Co-Fe-B bilayer, the effective fields with the Slonczewski-like torque (HSL) and field-like torque (HFL) symmetries were quantified by fitting [Fig. 2(a, b)]. Linear relation with the applied current density J [Fig. 2(c)] proves that these fields are induced by the current. These results would be important for further research on the current-induced magnetization switching using antiferromagnets. 
The authors acknowledge JSPS KAKENHI Grants 21K14522, 21K18189, the Core Research Cluster program, the MRAM program in CIES, Tohoku Univ.
Fig. 1: The second harmonic Hall resistance vs. field direction curves for Hall bar devices of RuO2/Co-Fe-B and Ru/Co-Fe-B.
Fig. 2: Fitting of HSL (a) and HFL (b). J dependence of μ0HSL and μ0HFL (c). 
**References:** [1] H. Honjo et al., IEDM Technical Digest 28.5, (2019). 
[2] M. Natsui et al., JSSC 56, 1116 (2020). 
[3] T. Tanaka et al., Phys. Rev. B 77, 165117 (2008) 
[4] A. Bose et al., Nature Electronics 5, 267 (2022). 
[5] H. Bai et al., Phys. Rev. Lett. 128, 197202 (2022). 

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Spin orbit torques in RuO2/Co

CoFeB/MgO-based p-MTJs are the de facto standard for current STT-MRAM [1-7]. A Co/Pt-based Ru-synthetic antiferromagnetic (SyF) coupling reference layer is widely used in p-MTJs to compensate the stray magnetic field from the reference layer owing to its high thermal stability [4]. However, to scale down MTJs, a higher stable reference layer is required [8]. To improve the reference layer stability, an Ir-SyF coupling layer has been proposed [9]. To utilize the Ir coupling layer in the Co/Pt-based reference layer, one needs to understand the relationship between MTJ properties and sputtering conditions. In this study, we have investigated the effect of the sputtering conditions on the magnetic and transport properties and also compared them with those of Ru. We prepared two types of stacks with Co/Pt-based Ir- and Ru-SyF coupling layers, as shown in Figs. 1-2. Ir was deposited by two sputtering conditions, denoted as Condition-A and Condition-B. The stacks were annealed at 400°C for 1 h. The magnetic and transport properties of the MTJ stacks were measured using VSM and CIPT method, respectively. The stack with Condition-A exhibits a larger exchange coupling field Hex than that with Condition-B (Fig.1). In the EDX line analysis (Fig.2), Ir profile for Condition-A is wider than that for Condition-B, resulting in lower Hex for Condition-B. A possible reason for the different intermixing degree of Ir and Co seems to be that the atoms for Condition-A recoiled from the Ir sputtering target have a higher energy than the atoms for Condition-A[10]. The stacks with the Ir-SyF show lower TMR ratio (Fig.2), resulting from the Pt diffusion into the MgO barrier. In contrast, the stacks with Ru show high TMR ratio because the Pt diffusion is suppressed. These results indicate that the suppression of Ir and Pt diffusion in the reference layer is crucial to achieving robust MTJs that can withstand annealing. This work was supported by CIES’s Industrial Affiliation on the STT MRAM program, CIES Consortium, JST-OPERA, and NEDO.

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Fig. 1 Coupling layer thickness dependence of Hex for SyF.

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Fig. 2 EDX depth profiles of the stacks with (a) Ir for Condition-B, (b) Ir for Condition-A , and (c) Ru.

Influence of sputtering conditions on the magnetic properties of Co/Pt-based iridium-synthetic antiferromagnetic coupling reference layer

Spin-orbit torque (SOT)-induced magnetization switching in heavy metal/ferromagnetic (HM/FM) heterostructures has attracted immense interest owing to its potential for the development of high-performance spintronics devices [1-4]. The SOT has been induced by two orthogonal components: the Slonczewski-like field (HSL) and the field-like field (HFL) which are equally capable of magnetization reversal [5,6]. Conventionally, HSL is considered to originate from the bulk spin Hall effect (SHE) induced by the strong spin-orbit coupling of bulk HM [1], and HFL is related to the interfacial Rashba Edelstein effect (REE) originating from the broken inversion symmetry [7]. While HFL was weaker than HSL in many HM-based SOT structures, obscuring the physics of HFL, a sizable value of HFL was reported in Ta-based SOT structures [8]. However, the detailed investigation of HFL has not been performed, and the further manipulation of HSL and HFL has not been reported. Herein, to investigate more details on the contribution of HSL and HFL to the SOT efficiency, we choose Ta/Co-Fe-B system and study the change in HSL and HFL by oxygen incorporation into Ta-layer. Furthermore, the SOT origin in these structures is also discussed. The in-plane magnetized Ta-O (tTa-O nm)/Co20Fe60B20 (1.2 nm)/MgO (1.3 nm)/Ta (1.0 nm) stacks with different Ta-O layer thickness (tTa-O) were fabricated on thermally oxidized Si substrates by RF sputtering. Ta-O underlayer (UL) was formed by in-situ natural oxidization of Ta layer at 0.3; 0.1; 0.03 Pa for 30 minutes. A stack with a 1 nm-Ta UL was also prepared as a reference sample for comparison. HSL and HFL were evaluated using an extended harmonic Hall measurement for Hall bar devices, which enables the exclusion of the thermoelectric effects. An in-plane external magnetic field with various strengths (Hext) was rotated in the film plane, whose direction was represented by the azimuthal angle θ. 10 Hz AC current was applied to the channel and the first and second harmonic Hall voltages were measured by a lock-in amplifier. Figures 1 (a) and (b) show the in-plane magnetic field angle θ dependence of the first (Rω) and second (R2ω) harmonic Hall resistance at different Hext. Rω component for the in-plane magnetized sample is dominated by the planar Hall resistance, following the equation: R1ω=RPHEsin2θ (1), where RPHE is the planar Hall resistance coefficient. From R2ω component, HSL and HFL can be obtained as follows: R2ω = - {RAHE[HSL/(Hext-Hk,eff)+RT]}×cosθ + [2RPHE(HFL+HOE)/Hext]×(2cos3θ-cosθ) (2), where RT is the second harmonic Hall resistance coefficient due to the thermoelectric voltage, HOE is the current-induced Oersted field. RAHE is the anomalous Hall effect (AHE) resistance and Hk,eff is the effective anisotropy field, evaluated from the AHE measurement in the perpendicular magnetic field direction. From the cosθ contribution of R2ω as a function of 1/(µ0Hext-µ0Hk,eff) (upper panel of Fig. 1 (c)); and the (2cos3θ - cosθ) contribution of R2ω as a function of 1/(µ0Hext) (lower panel of Fig. 1 (c)), HSL and HFL were evaluated. Figures 2 (a) and (b) show the oxygen pressure (POxygen) dependence of µ0HSL/I, µ0HFL/I, and HFL/HSL for samples with Ta-O UL of 1 nm, in comparison with the reference sample (dots at POxygen = 0 Pa). Herein, I is the total current applying to the device. The sign and magnitude of µ0HSL/I and µ0HFL/I are consistent with those in the previous work [8]. For all samples, the values of µ0HSL/I and µ0HFL/I are in the same order, implying their equal contribution to the total SOT in both Ta and Ta-O UL systems. As POxygen increases, the magnitude of µ0HSL(FL)/I decrease, which might be attributed to the low oxygen-incorporating in the system. However, HFL/HSL increases with the increase of POxygen meaning that this ratio would be enhanced by oxidization [9]. Figures 2 (c) and (d) show the tTa-O dependence of µ0HSL/I, µ0HFL/I, and HFL/HSL for stack with Ta-O UL oxidized at 0.3 Pa. The absolute values of µ0HSL/I, µ0HFL/I increase and converge to a certain value as tTa-O increases, suggesting the coexistence of bulk and interface effects originating from SHE, and REE [6]. The absolute value of µ0HFL/I is higher than that of µ0HSL/I suggesting the stronger interfacial REE-related origin than bulk SHE-related origin, which also can be observed in Fig. 2 (d) with the increase of HFL/HSL ratio as tTa-O increases. These results are of importance to understand the SOT induced magnetization switching, as well as to design the spintronic devices using oxidized HM/FM systems. The authors acknowledge the Core Research Cluster program of Tohoku University, Japan. This work was supported by JSPS KAKENHI Grant Number 19K23583, 19H00844, and by CSRN, and Cooperative Research Project from CSIS, Tohoku University, Japan. This research has been partly carried out at the Laboratory for Nanoelectronics and Spintronics, RIEC, Tohoku University. Fig. 1: Field direction dependence of Rω (a) and R2ω (b) at different magnetic field strength. The black curves are fitting data with eqs. (1), (2). (c) The cosθ contribution of R2ω as a function of 1/(μ0Hext - μ0Hk,eff) (upper); and the (2cos3θ - cosθ) contribution of R2ω as a function of 1/(μ0Hext) (lower). Fig. 2: Oxygen pressure dependence of μ0HSL(FL)/I (a) and HFL/HSL (b) for samples with 1 nm thick UL. tTa-O dependence of μ0HSL(FL)/I (c) and HFL/HSL (d) for samples with Ta-O UL oxidized at 0.3 Pa.

**References:**

[1] L. Liu et al., Science 336, 555, (2012). [2] S. Fukami et al., in Symp. VLSI Technol. Dig. Tech. Papers, pp. 60–61, (2016). [3] H. Honjo et al., IEDM Technical Digest 28.5, (2019). [4] M. Natsui et al., in Symp. VLSI Technol. Dig. Tech. Papers, CM2.2, (2020). [5] C. Zhang et al., Appl. Phys. Lett. 107, 012401, (2015). [6] Y. Du et al, Phys. Rev. Appl. 13, 054014, (2020). [7] V. M. Edelstein, Solid State Commum. 73, 233, (1990). [8] C. Zhang et al., Appl. Phys. Lett. 103, 262407, (2013). [9] H. An et al., Sci. Adv. 4, eaar2250, (2018).

**Images:**

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 Spin-orbit torque in naturally oxidized Ta-O/Co-Fe-B/Mg-O/Ta structures

Over the past decades, there has been a significant development in spin-transfer-torque magneto-resistive random-access memory (STT-MRAM). The development was brought by CoFeB/MgO-based magnetic tunnel junctions (MTJs) with interfacial perpendicular magnetic anisotropy (IPMA) [1]. Because thermal stability factor (Δ) is proportional to effective PMA constant and the volume of the free layer in a coherent magnetic rotation model [2] and proportional to square root of effective PMA constant and the diameter of the free layer in a domain propagation model [3], Δ decreases with the decrease of MTJ diameter. To achieve great enough Δ value in a smaller MTJ, increasing the number of CoFeB/MgO interfaces has been proposed to increase PMA [4-6]. In double interface MTJ with MgO/CoFeB/insertion layer/CoFeB/MgO free layer, thin non-magnetic layer such as Ta, W or Mo was used as the insertion layer to obtain higher PMA and tunnel magnetoresistance (TMR) ratio [4,7]. When the insertion layer is too thick, the two CoFeB layers in the free layer behave magnetically independent, resulting in no benefits on Δ despite the double interface because effective free-layer volume does not increase. To obtain a benefit of double CoFeB / MgO interfaces on Δ, a great enough magnetic coupling between the two CoFeB layers is required. It is important to clarify the effect of magnetic coupling (the energy constant of unit area: Jcpl) on Δ. However, it is not clear how the magnitude of Jcpl affects the increase in Δ, although we achieved a strong Jcpl in the double interface MTJs [8]. In this study, we propose a method to calculate critical Jcpl value to maximize Δ for free layers of MgO/CoFeB/insertion layer/CoFeB/MgO in the double interface MTJs with various MTJ diameters. We introduced magnetic coupling energy between two magnetic layers in the free layer based on magnetic domain propagation model (Model A) and magnetization coherent rotation model (Model B) as shown in Fig.1. When d was larger than the critical diameter dc (= 16/π [As /Keff ]1/2 ), Model A was applied, in which a blue-colored domain with downward magnetization nucleates from an edge and propagates in each FL1 and FL2. The domain nucleation and propagation behavior from edge region assumed in this model is consistent with the results of micromagnetic simulations of the MTJs with CoFeB / MgO interfaces [9,10]. The nucleated domain area Ω1 in FL1 and Ω2 in FL2 were decided by angles φ1 and φ2. A magnetic energy E (φ1, φ2) in Model A was calculated as follows. E (φ1, φ2) = Λ(φ1) t1 σw1 + Λ(φ2) t2 σw2 + 2Jcpl |Ω(φ1) - Ω(φ2)|---(1) σw1 = 4(As Keff1)1/2, σw2 = 4(As Keff2)1/2 ----(2) Λ(φi) = d (π/2–φi) tan φi, (i = 1, 2)---(3) Ω(φi) = d2 φi/4 + d2/4 { (π/2 – φi）tan2φi – tanφi }, (i = 1, 2) ---(4) The 1st and 2nd terms in the equation (1) express magnetic domain wall energies of FL1 and FL2, respectively. The third term expresses magnetic coupling energy between FL1 and FL2. Λ is wall length and σw is wall energy per unit area. As is stiffness constant of the free layer. When d is smaller than the critical diameter dc, coherent rotation model (Model B) was applied, in which E (φ1, φ2) was calculated as follows. The 1st and 2nd terms in the equation (5) express PMA energies of FL1 and FL2. The third term expresses magnetic coupling energy between FL1 and FL2. E (φ1, φ2) =πd2/4 { Keff1t1 (1–cos2φ1) + Keff2t2 (1–cos2φ2)+ Jcpl (1–cos(φ1–φ2）} ---(5) Fig. 2 shows Jcpl dependence on Δ for the MTJs with various diameters. We used Keff1=Keff2 = Keff_film +(1-Nz-x) Ms2/(2μ0). Keff_film = 2.25x105 J/m3, t1=t2=1.2nm, Ms = 0.97 T and As = 5.4 pJ/m, which were experimental values in the double interface MTJ [8]. The calculated critical diameter dc was 23 nm. With increasing Jcpl, Δ increases and saturates over a critical Jcpl value (Jcpl_c). The saturated Δ is twice as great as that of Jcpl = 0. When Jcpl ≤ Jcpl_c, energy barrier is minimized in the paths where the magnetizations of FL1 (MFL1) and FL2 (MFL2) separately reverse (not shown in Figures). On the other hand, when the Jcpl ≥ Jcpl_c, energy barrier is minimized in the path where MFL1 and MFL1 reverse simultaneously, resulting in maximum Δ. Jcpl_c value rapidly increases with the decrease of the MTJ diameter in Model A (d ≥ dc), and gradually increases in Model B (d ≤ dc). Jcpl_c value, which is equivalent with Keff1t1 ( =Keff_filmt1 +(1-Nz-x) Ms2/(2μ0)t1) when d ≤ dc, increases with the decrease of the demagnetic field coefficient Nz-x. We have reported that Jcpl of the blanket film of the double interface MTJ with a W insertion layer thickness of 0.2 nm reaches 0.44 mJ/m2 [8]. Therefore, all the Jcpl_c values obtained for various diameters based on these models are experimentally achievable. Consequently, we successfully established a method to predict critical Jcpl values to maximize Δ for double interface MTJ elements. This work is supported by CIES’s Industrial Affiliation on STT-MRAM program, CIES Consortium, JST-OPERA, and CAO-SIP.

**References:**

[1] S. Ikeda et al., Nat. Mater. 9, 721 (2010). [2] M. P. Sharrock, IEEE Trans. Magn. 26, 193 (1990). [3] G. D. Chaves-O’Flynn et al. Phys. Rev. Applied 4, 024010 (2015). [4] H. Sato et al. Appl. Phys. Lett. 101, 022414 (2012). [5] K. Nishioka et al., IEEE Trans. Electron Devices 67, 995 (2020). [6] S. Miura et al., IEEE Trans. Electron Devices 67, 5368 (2020). [7] J. H. Kim et al., Sci. Rep. 5, 16903 (2015). [8] K. Nishioka et al., MMM 2020 conference, O2-06. [9] I. Volvach et al., MMM 2020 conference, O2-07. [10] I. Volvach et al., Appl. Phys. Lett. 116, 192408 (2020).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/5ab35ffac43499d4ee18b9ab65311340.jpg)
Fig.1 Models to calculate magnetic energy surfaces of the free layer with double-interface MTJs. Model A is a magnetic domain propagation model and Model B is a magnetic coherent rotation model, introducing magnetic coupling energy between two FLs. Keff1t1 and Keff2t2 are effective perpendicular magnetic anisotropy const per unit area which MgO layers give rise to FL1 and FL2, respectively.

![](https://assets.underline.io/uploads/markdown_image/1/image/862f0cc74254d646eefaa20472e314b2.jpg)
Fig.2 Jcpl dependence on Δ for MTJs with various diameters. Jcpl_c is a Jcpl value which saturates Δ for a MTJ diameter d. (a) Model A is applied when d ≥ dc = 23 nm. (b) Model B is applied when d < dc = 23nm.

 Effect of magnetic coupling between two CoFeB layers on thermal stability in perpendicular magnetic tunnel junctions with MgO/CoFeB/insertion layer/CoFeB/MgO free layer

I. INTRODUCTION Ferromagnetic resonance (FMR) measurement technique is well known to be one of the important methods to evaluate the dynamic magnetic properties such as the effective saturation magnetization, damping constant, and so on, of magnetic thin film materials [1]. Recently, a broadband FMR (B-FMR) measurement technique with a coplanar waveguide (CPW) and a vector network analyzer (VNA) has been widely used to investigate them [2]-[4]. The reason for this is that the films are excited in a very wide range of RF frequency changing external magnetic field intensity and its high sensitivity allows for downsizing of the samples to the micro-scale, which provide detailed information on very thin films subjected to practical use under a variety of circumstances. In general, for FMR measurement technique, it is pointed out that the non-uniform magnetization precession mode occurs by the applied direction of external magnetic field. However, for the B-FMR measurement technique, this problem still remains an open question. Herein, we chose Fe-M binary alloys as magnetic thin film materials, investigated the in-plane and out-of-plane dynamic magnetic properties of Fe-M binary alloy thin films using the B-FMR measurement technique, and further discussed the difference between these properties in their films by the applied direction of external magnetic field in detail. II. EXPERIMENTAL PROCEDURES 10-nm thick Fe-M (M=Ni, Si, Co, and Ga) binary alloy thin films were fabricated by DC magnetron sputtering onto glass substrates. These film compositions were determined by energy dispersive x-ray spectroscopy (EDX), while the crystallographic structure of these films was performed using the high-angle XRD analysis and TEM observation, indicating that all film structures are polycrystalline and bcc phase. As for the dynamic magnetic properties of these films, the in-plane and out-of-plane effective saturation magnetization (4πMs,eff, //, 4πMs,eff,⊥) and the in-plane and out-of-plane effective damping constants (α//, α⊥) were evaluated by the B-FMR measurement technique. The external magnetic field was applied parallel to the CPW plane and swept between -0.01–0.3 T at various resonance frequencies. On the other hand, 4πMs,eff,⊥ and α⊥ of the films were determined in an external magnetic field applied perpendicular to the CPW plane swept between 1.4–3.3 T at various resonance frequencies. These measurements were performed at room temperature. The other magnetic properties of these films were already described in [4], [5] in detail. III. RESULTS AND DISCUSSION 4πMs,eff, //, 4πMs,eff,⊥, α//, and α⊥ of 10-nm thick Fe-M (M=Ni, Si, Co, and Ga) thin films are summarized in Table 1 and 2. As be noticed in Table 1, 4πMs,eff, // values of these films became higher or lower than those of saturation magnetization (4pMs) evaluated by VSM, which might be attributed to the appearance of negative or positive surface anisotropy. On the other hand, 4pMs,eff,⊥ of these films almost agreed with those of 4πMs evaluated by VSM, meaning that the easy magnetization direction of every film is in the film plane. As for the effective damping constant (Table 2), in case of M=Ni, α// value was almost equal to that of α⊥, indicating that the intrinsic damping can be obtained regardless of the external magnetic field direction. On the other hand, in case of M=Si, Co, and Ga, these damping depend on the external field direction, and these α// values were higher than those of α⊥. The reason for these differences might be that the non-uniform magnetization precession mode appears by magnetic inhomogeneities such as the two-magnon scattering, the anisotropy dispersion, and so on when the external magnetic field is applied in the film plane. On the basis of these results, it is suggested that the external magnetic field direction must be seriously selected by the kind of magnetic materials as their effective damping constants are evaluated by the B-FMR measurement technique. ACKNOWLEDGEMENT The authors thank Prof. Satoshi Okamoto, and Associate Prof. Nobuaki Kikuchi at Tohoku University for performing the photo lithography and measuring XRD diffraction patterns. This work was supported in part by KAKENHI (JP17H03226) from MEXT, Japan. This work was supported in part by CSIS, CSRN, and CIES, Tohoku University. This work was also supported in part by the ASRC in Japan.

**References:**

[1] J. H. E. Griffiths, Nature 148, 670 (1946)., [2] G. Counil et al., J. Appl. Phys. 95, 5646 (2004). [3] M. Toda et al., J. Magn. Soc. Jpn. 31, 435 (2007). [4] Y. Endo et al., J. Appl. Phys. 109, 07D336 (2011). [5] Y. Kawabe et al., Trans. Magn. Soc. Jpn. 3, 34 (2019).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/4fc57e3b82465a0fe4857ce200b27537.jpg)
Table 1. Comparison between the in-plane and out-of-plane effective saturation magnetization (4πMs,eff, //, 4πMs,eff,⊥) for Fe-M binary alloy thin films.

![](https://assets.underline.io/uploads/markdown_image/1/image/b17d7443a35e49fed8dce0227d6c370d.jpg)
Table 2. Comparison between the in-plane and out-of-plane effective damping constants (α// and α⊥) for Fe-M binary alloy thin films.

 Study on comparison between in-plane and out-of-plane dynamic magnetic properties for Fe-M binary alloy thin films

FOB: Spin Hall and Spin-Orbit Torque I

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Presentations

Spin orbit torques in RuO2/Co

Influence of sputtering conditions on the magnetic properties of Co/Pt-based iridium-synthetic antiferromagnetic coupling reference layer

Spin-orbit torque in naturally oxidized Ta-O/Co-Fe-B/Mg-O/Ta structures

Effect of magnetic coupling between two CoFeB layers on thermal stability in perpendicular magnetic tunnel junctions with MgO/CoFeB/insertion layer/CoFeB/MgO free layer

Study on comparison between in-plane and out-of-plane dynamic magnetic properties for Fe-M binary alloy thin films

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