France



Ability to control the spin orientation in antiferromagnets (AFMs) has recently attracted a lot of attention [1]. In our recent work we performed systematic studies on magnetic properties of NiO in Fe/NiO/Cr/MgO multilayers for various Cr thicknesses [2]. We proved existence of a continuous strain-induced spin reorientation transition (SRT) in NiO from nearly out-of-plane direction in Fe/NiO/MgO to the in-plane direction in Fe/NiO/Cr. For ultrathin NiO layers the in plane direction of AFMs spins was determined to be perpendicular to Fe spins due to spin-flop interaction.&lt;br/&gt;In this contribution we studied magnetic properties of NiO in NiO/MgO(&lt;i&gt;d&lt;sub&gt;MgO&lt;/sub&gt;)&lt;/i&gt;/Cr(20 nm)/MgO(001) for different MgO thicknesses, &lt;i&gt;d&lt;sub&gt;MgO&lt;/sub&gt;&lt;/i&gt;. We also investigated how the adjacent ferromagnetic coverage layer influence the spin structure in NiO.&lt;br/&gt;X-ray Magnetic Linear Dichroism (XMLD) technique was used to determine magnetic ordering in NiO. The XAS spectra were collected using a linearly polarized x-ray beam for the Ni&lt;sup&gt;2+ &lt;/sup&gt;L&lt;sub&gt;2&lt;/sub&gt; edge at normal (γ = 0°) and grazing (γ = 60°) x-ray incident angles. The XAS spectra were measured as a function of MgO buffer layer thickness for NiO/MgO/Cr, MgO/NiO/MgO/Cr and Fe/NiO/MgO/Cr. Figure 1 shows L&lt;sub&gt;2&lt;/sub&gt; ratio difference, defined as ΔRL&lt;sub&gt;2&lt;/sub&gt;= RL&lt;sub&gt;2&lt;/sub&gt;(γ = 0°) – RL&lt;sub&gt;2&lt;/sub&gt;(γ = 60°), as a function of &lt;i&gt;d&lt;sub&gt;MgO&lt;/sub&gt;&lt;/i&gt; for NiO with different capping layers. For both, capped and non-capped NiO layers we noted an increase in ΔRL&lt;sub&gt;2&lt;/sub&gt; with an increase in &lt;i&gt;d&lt;sub&gt;MgO&lt;/sub&gt;&lt;/i&gt;. According to the previous studies [3] increase of ΔRL&lt;sub&gt;2&lt;/sub&gt; can be interpreted as enhancement of the out-of-plane component of NiO spins. Thus, our results show that insertion of MgO layer between NiO and Cr leads to development of out-of-plane NiO spin direction, which can be understand in terms of tensile strain exerted by MgO buffer layer on NiO.&lt;br/&gt;Interestingly, while there are no significant differences between ΔRL&lt;sub&gt;2&lt;/sub&gt;(&lt;i&gt;d&lt;sub&gt;MgO&lt;/sub&gt;&lt;/i&gt;) dependences for MgO-capped and uncapped NiO layers (Fig. 1, circles, triangles), a change of ΔRL&lt;sub&gt;2&lt;/sub&gt; as a function of MgO thickness is much more pronounced for Fe/NiO/MgO/Cr (Fig. 1, stars). Moreover, only for Fe/NiO/MgO/Cr ΔRL&lt;sub&gt;2 &lt;/sub&gt;changes a sign with an increase of &lt;i&gt;d&lt;sub&gt;MgO&lt;/sub&gt;&lt;/i&gt; which indicates that NiO spins undergo an SRT from in-plane to out-of-plane direction. To investigate the influence of Fe/NiO interactions on SRT in NiO a dedicated sample was prepared. Wedge-shaped Fe layer was grown on NiO(2 nm)/MgO(8.5 nm)/Cr(20 nm)/MgO(001), where the &lt;i&gt;d&lt;sub&gt;MgO&lt;/sub&gt;&lt;/i&gt; layer was chosen to be near the critical thickness of MgO for which SRT in NiO was observed. To elucidate if modulation of magnetic structure in NiO is influenced by magnetic properties of Fe or rather by strains from capping layer, half of the sample was capped with wedge-shaped Cr. Due to the small lattice mismatch between Cr and Fe (1.5%) strain-induced changes in magnetic properties of NiO should be similar for Cr/NiO/MgO and Fe/NiO/MgO. Prior to XAS measurement the sample was magnetized along the Fe[001] (NiO[110]) in-plane direction (φ = 0°) to study the influence of Fe magnetization on NiO spin structure. Figure 2(a) shows ΔRL&lt;sub&gt;2&lt;/sub&gt; dependence on &lt;i&gt;d&lt;sub&gt;Fe(Cr)&lt;/sub&gt; &lt;/i&gt;for two azimuthal angles, φ. In case of Cr/NiO, for both φ = 0° and φ = 90° we noted no considerable ΔRL&lt;sub&gt;2 &lt;/sub&gt;changes as a function of Cr thickness (&lt;i&gt;d&lt;sub&gt;Cr&lt;/sub&gt;&lt;/i&gt;) (Fig. 2(a) circles, triangles). For Fe/NiO we noted a drastic difference in ΔRL&lt;sub&gt;2&lt;/sub&gt;(&lt;i&gt;d&lt;sub&gt;Fe&lt;/sub&gt;&lt;/i&gt;) dependence for φ = 0° and φ = 90°. Strong ΔRL&lt;sub&gt;2&lt;/sub&gt; dependence on &lt;i&gt;d&lt;sub&gt;Fe&lt;/sub&gt;&lt;/i&gt; observed for φ = 90° and a lack of change in ΔRL&lt;sub&gt;2&lt;/sub&gt; as a function of &lt;i&gt;d&lt;sub&gt;Fe&lt;/sub&gt;&lt;/i&gt; for φ = 0° was noted for Fe/NiO/MgO/Cr, which brings two conclusions. First, a change of sign in ΔRL&lt;sub&gt;2&lt;/sub&gt;(&lt;i&gt;d&lt;sub&gt;Fe&lt;/sub&gt;&lt;/i&gt;) obtained for φ = 90° proves that SRT from in-plane to out-of-plane direction in NiO can be obtained as a function of Fe thickness. Dependence of ΔRL&lt;sub&gt;2&lt;/sub&gt;(&lt;i&gt;d&lt;sub&gt;Fe&lt;/sub&gt;&lt;/i&gt;) on azimuthal angle in Fe/NiO/MgO together with a lack of ΔRL&lt;sub&gt;2&lt;/sub&gt; dependence on &lt;i&gt;d&lt;sub&gt;Cr&lt;/sub&gt;&lt;/i&gt; in Cr/NiO/MgO prove that reorientation of NiO spins is not related to the strains exerted by capping layer, but to the magnetism of Fe layer. Second, a difference between ΔRL&lt;sub&gt;2&lt;/sub&gt;(&lt;i&gt;d&lt;sub&gt;Fe&lt;/sub&gt;&lt;/i&gt;) for φ = 0° and φ = 90° indicates that together with an increase in Fe thickness an in-plane anisotropy of NiO is developed.&lt;br/&gt;This conclusion was confirmed in XAS measurements collected as a function of φ (γ was set to 0) for different Fe thicknesses (Fig. 2(b)). While for &lt;i&gt;d&lt;sub&gt;Fe&lt;/sub&gt;&lt;/i&gt;= 2 Å no change in the RL&lt;sub&gt;2&lt;/sub&gt;(φ) was observed, for &lt;i&gt;d&lt;sub&gt;F&lt;/sub&gt;&lt;/i&gt;&lt;sub&gt;e&lt;/sub&gt;≥8 Å we noted well-pronounced extrema in the RL&lt;sub&gt;2&lt;/sub&gt;(φ) dependence for the NiO[110] and [1–10]. According to the previous studies [3], if the RL&lt;sub&gt;2&lt;/sub&gt; reaches an extremum when the polarization vector is parallel to the [110] NiO direction, the AFM spin alignment is determined from the maximum of the RL&lt;sub&gt;2&lt;/sub&gt;. Thus, an in-plane component of NiO spins is perpendicular to the spins of Fe.&lt;br/&gt;Our results show that spin structure of antiferromagnet can be modulated by strain induced by the buffer layer and emphasize the role of adjacent ferromagnetic layer on the magnetic structure in AFM.&lt;br/&gt;&lt;br/&gt;&lt;i&gt;This work was supported by the “Antiferromagnetic proximity effect and development of epitaxial bimetallic antiferromagnets – two routes towards next-generation spintronics” project, which is carried out within the Homing programme of the Foundation for Polish Science co-financed by the European Union under the European Regional Development Fund.&lt;/i&gt;

**References:**

[1] C. Song, Y.You and X.Chen, &lt;i&gt;Nanotechnology&lt;/i&gt;, Vol.29, p.112001(2018)&lt;br/&gt;[2] A. Koziol-Rachwal, M. Slezak, M. Zajac, &lt;i&gt;APL Mater.&lt;/i&gt;, Vol.8, p.061107 (2020)&lt;br/&gt;[3] W. Kim, E. Jin, J. Wu, &lt;i&gt;Phys. Rev. &lt;/i&gt;&lt;i&gt;B.,&lt;/i&gt; Vol.81, p.174416 (2010)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/3eb02b9b080b4178806eb96c5dd60035.jpg)
&lt;i&gt;Fig. &lt;/i&gt;&lt;i&gt;1&lt;/i&gt;&lt;i&gt;. XLD magnitudes of NiO in the NiO/MgO/Cr, MgO/NiO/MgO/Cr, Fe/NiO/MgO/Cr.&lt;/i&gt;

![](https://assets.underline.io/uploads/markdown_image/1/image/b44441a561c3ec343508212ae55e0015.jpg)
&lt;i&gt;Fig. 2. a) XLD magnitudes of NiO in the Fe/NiO and Cr/NiO as a function of d&lt;sub&gt;Fe(Cr), &lt;/sub&gt;b) Ni RL&lt;sub&gt;2 &lt;/sub&gt;(φ) as a function of polarization angle φ for Fe(2 Å)/NiO ,Fe(8 Å)/NiO, Fe(18 Å)/NiO and Fe(30 Å)NiO.&lt;/i&gt;

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 Strain and ferromagnetic proximity induced spin reorientation transition in NiO

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Ability to control the spin orientation in antiferromagnets (AFMs) has recently attracted a lot of attention [1]. In our recent work we performed systematic studies on magnetic properties of NiO in Fe/NiO/Cr/MgO multilayers for various Cr thicknesses [2]. We proved existence of a continuous strain-induced spin reorientation transition (SRT) in NiO from nearly out-of-plane direction in Fe/NiO/MgO to the in-plane direction in Fe/NiO/Cr. For ultrathin NiO layers the in plane direction of AFMs spins was determined to be perpendicular to Fe spins due to spin-flop interaction. In this contribution we studied magnetic properties of NiO in NiO/MgO(dMgO)/Cr(20 nm)/MgO(001) for different MgO thicknesses, dMgO. We also investigated how the adjacent ferromagnetic coverage layer influence the spin structure in NiO. X-ray Magnetic Linear Dichroism (XMLD) technique was used to determine magnetic ordering in NiO. The XAS spectra were collected using a linearly polarized x-ray beam for the Ni2+ L2 edge at normal (γ = 0°) and grazing (γ = 60°) x-ray incident angles. The XAS spectra were measured as a function of MgO buffer layer thickness for NiO/MgO/Cr, MgO/NiO/MgO/Cr and Fe/NiO/MgO/Cr. Figure 1 shows L2 ratio difference, defined as ΔRL2= RL2(γ = 0°) – RL2(γ = 60°), as a function of dMgO for NiO with different capping layers. For both, capped and non-capped NiO layers we noted an increase in ΔRL2 with an increase in dMgO. According to the previous studies [3] increase of ΔRL2 can be interpreted as enhancement of the out-of-plane component of NiO spins. Thus, our results show that insertion of MgO layer between NiO and Cr leads to development of out-of-plane NiO spin direction, which can be understand in terms of tensile strain exerted by MgO buffer layer on NiO. Interestingly, while there are no significant differences between ΔRL2(dMgO) dependences for MgO-capped and uncapped NiO layers (Fig. 1, circles, triangles), a change of ΔRL2 as a function of MgO thickness is much more pronounced for Fe/NiO/MgO/Cr (Fig. 1, stars). Moreover, only for Fe/NiO/MgO/Cr ΔRL2 changes a sign with an increase of dMgO which indicates that NiO spins undergo an SRT from in-plane to out-of-plane direction. To investigate the influence of Fe/NiO interactions on SRT in NiO a dedicated sample was prepared. Wedge-shaped Fe layer was grown on NiO(2 nm)/MgO(8.5 nm)/Cr(20 nm)/MgO(001), where the dMgO layer was chosen to be near the critical thickness of MgO for which SRT in NiO was observed. To elucidate if modulation of magnetic structure in NiO is influenced by magnetic properties of Fe or rather by strains from capping layer, half of the sample was capped with wedge-shaped Cr. Due to the small lattice mismatch between Cr and Fe (1.5%) strain-induced changes in magnetic properties of NiO should be similar for Cr/NiO/MgO and Fe/NiO/MgO. Prior to XAS measurement the sample was magnetized along the Fe[001] (NiO[110]) in-plane direction (φ = 0°) to study the influence of Fe magnetization on NiO spin structure. Figure 2(a) shows ΔRL2 dependence on dFe(Cr) for two azimuthal angles, φ. In case of Cr/NiO, for both φ = 0° and φ = 90° we noted no considerable ΔRL2 changes as a function of Cr thickness (dCr) (Fig. 2(a) circles, triangles). For Fe/NiO we noted a drastic difference in ΔRL2(dFe) dependence for φ = 0° and φ = 90°. Strong ΔRL2 dependence on dFe observed for φ = 90° and a lack of change in ΔRL2 as a function of dFe for φ = 0° was noted for Fe/NiO/MgO/Cr, which brings two conclusions. First, a change of sign in ΔRL2(dFe) obtained for φ = 90° proves that SRT from in-plane to out-of-plane direction in NiO can be obtained as a function of Fe thickness. Dependence of ΔRL2(dFe) on azimuthal angle in Fe/NiO/MgO together with a lack of ΔRL2 dependence on dCr in Cr/NiO/MgO prove that reorientation of NiO spins is not related to the strains exerted by capping layer, but to the magnetism of Fe layer. Second, a difference between ΔRL2(dFe) for φ = 0° and φ = 90° indicates that together with an increase in Fe thickness an in-plane anisotropy of NiO is developed. This conclusion was confirmed in XAS measurements collected as a function of φ (γ was set to 0) for different Fe thicknesses (Fig. 2(b)). While for dFe= 2 Å no change in the RL2(φ) was observed, for dFe≥8 Å we noted well-pronounced extrema in the RL2(φ) dependence for the NiO[110] and [1–10]. According to the previous studies [3], if the RL2 reaches an extremum when the polarization vector is parallel to the [110] NiO direction, the AFM spin alignment is determined from the maximum of the RL2. Thus, an in-plane component of NiO spins is perpendicular to the spins of Fe. Our results show that spin structure of antiferromagnet can be modulated by strain induced by the buffer layer and emphasize the role of adjacent ferromagnetic layer on the magnetic structure in AFM. This work was supported by the “Antiferromagnetic proximity effect and development of epitaxial bimetallic antiferromagnets – two routes towards next-generation spintronics” project, which is carried out within the Homing programme of the Foundation for Polish Science co-financed by the European Union under the European Regional Development Fund.

**References:**

[1] C. Song, Y.You and X.Chen, Nanotechnology, Vol.29, p.112001(2018) [2] A. Koziol-Rachwal, M. Slezak, M. Zajac, APL Mater., Vol.8, p.061107 (2020) [3] W. Kim, E. Jin, J. Wu, Phys. Rev. B., Vol.81, p.174416 (2010)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/3eb02b9b080b4178806eb96c5dd60035.jpg)
Fig. 1. XLD magnitudes of NiO in the NiO/MgO/Cr, MgO/NiO/MgO/Cr, Fe/NiO/MgO/Cr.

![](https://assets.underline.io/uploads/markdown_image/1/image/b44441a561c3ec343508212ae55e0015.jpg)
Fig. 2. a) XLD magnitudes of NiO in the Fe/NiO and Cr/NiO as a function of dFe(Cr), b) Ni RL2 (φ) as a function of polarization angle φ for Fe(2 Å)/NiO ,Fe(8 Å)/NiO, Fe(18 Å)/NiO and Fe(30 Å)NiO.

technical paper

Introduction: Exchange Bias(EB) originates from the interfacial interaction between a ferromagnetic (FM) and an Antiferromagnetic (AFM) material, after cooling the sample in a magnetic field below the Neel temperature (TN) of the AFM [1]. The main feature of EB is training effect which is a gradual reduction of EB field when hysteresis loop (n) measured many times at constant temperature. The training effect is inversely dependence on FM layer thickness and interface structure, etc [2]. We have carried out the Kerr microscopy measurements to get direct information about magnetization irreversible of FM layers along with hysteresis loops. Experimental details: Thin films of Al(10nm)/Fe(30nm)/Fe57(10nm)/Fe(30nm)/Si substrate were deposited using Ion-beam sputtering technique. The films were implanted with 35 keV Fluorine ions. With this ion energy, the maximum energy has been deposited only at the centre of the Fe film to obtain FeF2 at the center, which is an antiferromagnet and its Neel temperature is 78 K. For magnetic characterization, the measurements were done using in-plane Magneto-optic Kerr effect microscope (MOKE) at 50 K after field cooling from room temperature. Results and discussions Fig. 1(a) shows the hysteresis loop as function of number of loops. It clearly shows the exchange bias and it becomes quite constant. Similar observation has also been showed for the sample irradiated with 5×1016 F-ions/cm2. Magnetic domain evolution during the magnetization irreversible and the training effect have been studied after F-ion implantation at 50 K. Kerr microscopy measurements were carried out at different temperatures after the field-cooling across the Neel temperature of the antiferromagnetic (AFM) FeF2 layer. It is observed that with the increasing anti-ferromagnetic (AFM) FeF2 layer thickness, the exchange bias increase and the training effect becomes weaker. Analysis of the temperature variation of the exchange bias field indicates the disorder decreases at the FM-AFM interface with the increasing F-Ion fluences. Higher fluence specimen indicates a behavior that is changing from non-equilibrium state to equilibrium state and behavior of training effect is suggesting step-like. For 1×1017 ions/cm2 sample, change in the exchange bias field after measuring the 6th cycle is 4.42% while in case of 5×1016ions/cm2 27.32%. Remarkably, for the largest fluences, the relative training becomes smaller than the one for the lower fluence. It indicates higher ion fluence EB system is suitable for device application.

**References:**

 [1] W. H. Meiklejohn and C. P. Bean, Phys. Rev. 102, 1413(1956). [2] M. Gruyters and D. Riegel, Phys. Rev. B 63, 052401 (2000).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/a97c3e88eac11173ccbf237d54b463aa.jpg)
Figure 1: (a) The hysteresis loops as a function of number of loops (n) for the specimen 1×1017 ions/cm2 (b) Value of Exchange bias for both the films implanted with ion fluences 5×1016 ions/cm2 and 1×1017 ions/cm2 at 50 K.

 Observation of training effect in Fe thin film implanted with F+ ions

Magnetic tunnel junctions (MTJs) with perpendicular magnetic anisotropy (PMA) have advantages of high packing density due to their high thermal stability, efficient current-induced magnetic switching, and unrestricted cell aspect ratio.[1] Therefore, PMA-based MTJs (p-MTJs) have broad application prospects on high-density and non-volatile spin-transfer torque magnetic random access memories (STT-MRAMs) and other spintronics devices.[2-4] In particular, p-MTJs with CoFeB/MgO structures show high tunnel magnetoresistance ratio.[4] In the W/CoFeB/MgO stacks, the interfacial magnetic anisotropy (Ki) of 1.98 mJ/m2 is acquired after annealing at 350 °C.[5] To further improve Ki, it is an effective way to use MgO/CoFeB/X/CoFeB/MgO double-interface free layer structures, where the X represents the spacer material. [6] However, higher Ki is needed to scale down the MTJs devices and maintain high thermal stability.[7] In this work [8], we study the magnetic anisotropy of the double-interface free layer stacks with different annealing temperatures and different spacer materials. The samples of Ta(5)/MgO(2)/Co20Fe60B20(t)/X(0.4) /Co20Fe60B20(1.2)/MgO(2)/Ta(3) are deposited by sputtering system, where the numbers in parentheses indicate the thickness in nanometers, X including Ta, W, and Mo. The thickness of the bottom CoFeB layer t ranges from 0.4~3.2 nm. The films with Mo spacer layer exhibit the highest Ki of 4.06 mJ/m2 compared with the Ta- and W-based films. Furthermore, the Ki of Mo-based films remains at 2.93 mJ/m2 even after annealing at 500 °C for 1 h. Fig.1a and 1b show the structures with bottom CoFeB thickness of 0.8 nm and with spacer materials of Mo and W. As shown in the M-H (Magnetization M versus magnetic field H) curves, all the films show good PMA. Fig.1c and 1d show the corresponding high-resolution transmission electron microscopy (HR-TEM) and energy-dispersive X-ray spectroscopy (EDX) images. Saturation magnetization Ms and the total thickness of magnetic dead layer tDL can be extracted from the curves ( M/A versus tCoFeB curves), where A is the area of the sample and tCoFeB is the total thickness of CoFeB layers. The Ki can be extracted from the intercept of the Keff*teff versus teff curves, where Keff is effective magnetic anisotropy energy density and teff = tCoFeB- tDL. As shown in Fig.2a and 2b, Ki of 2.15±0.11 mJ/m2 and 2.90±0.34 mJ/m2 can be obtained for Ta- and W-based stacks, respectively. More importantly, the maximum of 4.06±0.14 mJ/m2 is acquired for the Mo-based structure after annealing at 350 °C as shown in Fig.2c. According to previous study, when the required thermal stability factor is 70, if the Ki reached 4 mJ/m2 , the device size can be reduced to 10 nm.[8] It can be shown in Fig.2d- 2f that the Ki decreases with the increase of annealing temperature. Note that the Ki reaches 2.93 mJ/m2 even after annealing at 500 °C for 1 h, which demonstrates the high thermal stability of this structure. Revealed by the HR-TEM and EDX images, the sharper interface, weaker interdiffusion, better interfacial uniformity and higher crystallinity can be seen in the Mo-based stack than the W-based stack. These results show that Mo plays an important role in improving the film quality. When the annealing temperature increases, the interfacial mixing intensifies, which leads to the decrease of Ki in Mo-based films. In summary, we study the dependence of magnetic anisotropy and film quality of the MgO/CoFeB/X/CoFeB/MgO double-interface free layer structures on the spacer material and the annealing temperature. We find that, the Mo-based structure exhibits the strongest PMA (4.06 mJ/m2). Furthermore, Ki in the Mo-based stack is still as strong as 2.93 mJ/m2 even after annealing at 500 °C. The HR-TEM and EDX images of W- and Mo-based stacks show that the high thermal stability can be attributed to the high film quality in the Mo-based films. Our work provides a practical solution to scale down MRAM device while maintaining a high thermal stability.

**References:**

[1] Ikeda, S., Miura, K., Ohno, H. Nature materials, 9(9), 721-724 (2010). [2] Wang, M., Cai, W., Zhao, W. Nature electronics, 1(11), 582-588 (2018). [3] Peng, S., Zhu, D., Zhao, W. Nature Electronics, 1-8 (2020). [4] Wang, M., Cai, W., Zhao, W. Nature communications, 9(1), 1-7 (2018). [5] An, G. G., Lee, J. B., Hong, J. P. Acta Materialia, 87, 259-265 (2015). [6] Hwang, C. S. Advanced Electronic Materials, 1(6), 1400056 (2015). [7] Peng, S., Kang, W., Zhao, W. IEEE Magnetics Letters, 8, 1-5 (2017). [8] Cheng, H., Chen, J., Zhao, W. Advanced Electronic Materials, 6(8), 2000271 (2020).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/c74384cddd3ce3d115a89232da765fba.jpg)
Fig.1. (a,b) The schematic of Ta(5)/MgO(2)/Co20Fe60B20(0.8)/X(0.4) /Co20Fe60B20(1.2)/MgO(2)/Ta(3) with spacer material of (a) X=Mo and (b) X=W. (c,d) The corresponding HR-TEM images and EDX line profiles with spacer material of (c) X=Mo and (d) X=W. (e,f) The corresponding hysteresis loops with the in-plane (IP) field (red line) and out-of-plane (OP) field (black line) for different spacer materials of (e) X=Mo and (f) X=W. In the inset of (e) and (f), a M-H image under the small OP field with spacer material of (e) X=Mo and (f) X=W.

![](https://assets.underline.io/uploads/markdown_image/1/image/75ebcd2b300cee56aef34e4a13c0417a.jpg)
Fig.2. (a-c) The Ki fittings for MgO (2)/Co20Fe60B20(t )/X(0.4)/ Co20Fe60B20(1.2)/MgO(2) stacks under annealing at 350 °C with spacer layer X of (a) Ta, (b) W, and (c) Mo. (d-f) The Ki fittings for MgO(2)/Co20Fe60B20(t ) /Mo(0.4)/Co20Fe60B20(1.2)/MgO(2) stacks with the annealing temperature of (d) 400 °C, (e) 450 °C, and (f) 500 °C.

 Giant perpendicular magnetic anisotropy in Mo-based double-interface free layer structure for advanced magnetic tunnel junctions

Soft magnetic films consisting of CoFeB alloys are widely applied in magnetic sensor [1] and MRAM devices [2], but also used as cores for microfabricated transformers used to power specific subunits on microprocessors [3]. These applications typically require a well-defined in-plane uniaxial anisotropy axis which aligns the magnetic moments and permits operation modes where a field is applied either along the anisotropy axis to obtain a domain wall motion easy axis or a magnetization rotation governed magnetization process. Several implementations have been discussed in literature to obtain a well-defined uniaxial in-plane anisotropy axis and to tune the anisotropy. Apart from the geometry of a sensor element which can be chosen to obtain a shape anisotropy , oblique sputtering of seed layers for the successively deposited magnetic layer(s) [4], sputter deposition or post-deposition annealing in a magnetic field [5] have been employed. Here, we deposit CoFeB layers on Ta seed layers grown by oblique sputtering and compare the resulting anisotropies to those obtained by deposition of the magnetic layer in an applied field. Moreover, for the CoFeB layers on the oblique-sputtered seeds we compare the anisotropies obtained for different CoFeB-thicknesses to distinguish between a surface or interface induced anisotropy or a bulk anisotropy term. Specifically, we deposited Ta(2nm)/CoFeB(6nm)/Ta(Xnm)-seed films onto a Si-wafer covered by its native oxide using angles of 40°, 50° and 60° between the sample normal and the Ta source, and Ta-seed layer thicknesses from 2nm to 10nm. For an angle of 60° the Ta-seed layer was additionally deposited with a thickness of 15nm and 20nm. In addition, we deposited 4nm, 8nm, 16nm and 32nm CoFeB films on 4nm-thick Ta-seed layer deposited at an angle of 50°. Selected systems were deposited in fields of 4mT and 40mT, applied parallel to the substrate surface. VSM was then used to obtain the magnetization of the CoFeB films. Hard- and easy axis magnetization loops were measured with a Kerr Microscope to determine the in-plane anisotropy (see figure 1 & figure 2). We find that the in-plane anisotropy increases with the deposition angle and also the seed layer thickness. For a deposition angle of 60° the anisotropy increases from 4kJ/m3 to 25kJ/m3 with thicknesses from 2nm to 20nm. For smaller deposition angles, we measure a much smaller increase of in-plane anisotropy. Concerning the scientific question whether the anisotropy arises from a surface or a bulk effect, our analysis of CoFeB films with thicknesses from 4nm to 32nm revealed that the anisotropies stemming from the oblique-sputtered Ta-seed layers is well modeled by a surface anisotropy term. While with oblique sputtering, a wide range of in-plane anisotropies can be covered, the anisotropies remain much smaller when films are deposited in fields. However, it has already been demonstrated that with CoSm-alloys much larger anisotropies have been obtained [5]. To correlate the seed-layer morphology with the in-plane anisotropy, high resolution AFM imaging on the Ta-seed layer was performed. Our work shows that in-plane anisotropies in magnetically soft systems can be tuned over a wide range by oblique seed-layer depositions. However, the anisotropy arises from an interface effect and thus decays as 1/t with CoFeB film layers thickness "t". The anisotropy of CoFeB films deposited in an applied field does not depend on film thickness and remains rather small. We propose that alloying Sm into the CoFeB films may increase the field-induced anisotropy, while keeping the soft magnetic properties of the CoFeB films.

**References:**

[1] T. Ogasawara, M. Oogane and M. Al-Mahdawi; Sci. Rep. 9, 17018 (2019) [2] S. Zhanga, Y. Sua and X. Li; Appl. Phys. Lett. 114, 042401 (2019) [3] J. Michel et al. IEEE Transactions on Magnetics 55, 7, pp. 1-7, (2019) [4] R. D. McMichael, C. G. Lee and J. E. Bonevich; Journal of Applied Physics 88, 5296 (2000) [5] F. Magnus, R. Moubah and A. H. Roos; Appl. Phys. Lett. 102, 162402 (2013)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/6f342ce0ad4cb0acb37eff2b59d57a63.jpg)
Hard- and easy axis magnetization loops measured with Kerr Microscope of deposited Ta(2nm)/CoFeB(6nm)/Ta(Xnm)-seed layer at an angle of 60°

![](https://assets.underline.io/uploads/markdown_image/1/image/4484fc484b9b980d8019b0d9f4b39d69.jpg)
In-plane anisotropy of deposited Ta(2nm)/CoFeB(6nm)/Ta(Xnm)-seed layer at an angle of 60°

 Tuning the inplane anisotropy of CoFeB films

In this work, we present the results of polar magneto-optical Kerr effect (PMOKE) magnetometry and microscopy study of Ta/CoFeB/MgO stacks with a prototypical perpendicular easy axis system in spintronics applications [1]. Multilayer films with a structure of Ta (5 nm)/(Co0.25Fe0.75)75B25 (d=1.24÷1.60 nm)/MgO (2 nm)/Ta (2 nm) were deposited by dc/rf magnetron sputtering on thermally-oxidized Si substrates and after were annealed at 300 oC for 1 h under a perpendicular magnetic field 0.4 T in vacuum [2,3]. The magnetization processes statics, magnetization reversal dynamics and domain structures evolution driven by magnetic field were studied as a function of CoFeB thickness d. Magnetization curves are measured for samples with various d applying PMOKE magnetometry. For d=1.24 to 1.36 nm square-shaped hysteresis loops with full remanence, indicating an out-of-plane magnetization easy axis, are observed. For d=1.40 nm the shape of magnetization curve corresponds to the transitional state (at the edge of out-of-plane and in-plane magnetization states). For d =1.44 nm the magnetization curves demonstrate no hysteresis, which indicates on the reversible magnetization rotation and in-plane magnetization. The values of magnetic anisotropy parameters were deduced from the fitting of hard axis hysteresis curves using macrospin approximation model. In the case of the samples without in-plane magnetization additional 0.05 T magnetic field was applied to prevent domain structure formation. In the case of square hysteresis loop (d=1.24÷1.36 nm) magnetization reversal dynamics was studied exploiting PMOKE microscopy imaging technique driven by sequential magnetic field pulses (duration 200 ms) with a constant amplitude. Magneto-optical images were acquired in zero field after each pulse. The derived normalized magnetization m can be obtained as a function of number of field pulses expressed as time t using image processing techniques [4]. The dependence of magnetization m for direct process (starting from full saturated state) on time t for the sample with d= 1.24 nm is shown in Fig. 1. The logarithm of time t1/2, time when the sample has magnetization m=0, depends linearly on amplitude of pulses, which indicates on thermal activation of jump-like domain wall motion [5] and allows to calculate Barkhausen length lB » 120 nm [6]. Magnetization reversal in the stacks with out-of-plane anisotropy (d = 1.24÷1.36 nm) goes through domain nucleation and domain wall propagation. In the case of the samples with d = 1.24, 1.28 nm, while increasing H few nucleation point appears initially in the field of view and later quickly expands covering all the area excluding narrow domain between neighboring expanding domains named as narrow stripe domains (NSDs). NSDs are stabilized due to magnetostatic repulsion between there opposite domain walls and finite value of coercivity [7]. The increase of thickness d lead to significant reduction of straights cuts of NSDs, and for d = 1.32, 1.36 nm (close to dSRT) NSDs network is converted to the more complex dendrite like structures. The evolution of NSDs network in external out-of-plane magnetic field was studied exploiting PMOKE microscopy imaging technique driven by sequential magnetic field pulses (duration 500 ms) with a growing amplitude. The digital image processing of acquired images yielded normalized magnetization and NSDs length in the field of view. The obtained dependences of normalized magnetization and NSDs length for the stack with d=1.24 nm are presented in Fig. 2 as the functions of the amplitude of applied field pulses. NSDs appear after beginning of nucleation process at 1.25 mT, quickly reach maximum at 1.45 mT (when normalized magnetization mL » 0.92), then gradually annihilate and disappear after 5.5 mT. So saturation field is few times bigger than coercivity due to the NSDs. In conclusion, we have preformed an experimental study of the static and dynamic magnetization processes in ultrathin FeCoB films depending on their thickness. We find that magnetization reversal is occurred via thermal activated domain wall motion and is influenced by NSDs. The dependence of NSDs network length on applied out-of-plane magnetic field was studied.

**References:**

[1] S. Ikeda, K. Miura, H. Yamamoto, K. Mizunuma, H. D. Gan, M. Endo, S. Kanai, J. Hayakawa, F. Matsukura and H. Ohno, A perpendicular-anisotropy CoFeB–MgO magnetic tunnel junction,, Nature Mater. 9, 721–724 (2010), 10.1038/nmat2804 [2] S. Fukami, T. Suzuki, Y. Nakatani, N. Ishiwata, M. Yamanouchi, S. Ikeda, S. Kasai, and H. Ohno, Current-induced domain wall motion in perpendicularly magnetized CoFeB nanowire, Appl. Phys. Lett. 98, 082504 (2011), 10.1063/1.3558917 [3] C. Burrowes, N. Vernier, J.-P. Adam, L. Herrera Diez, K. Garcia, I. Barisic, G. Agnus, S. Eimer, Joo-Von Kim, T. Devolder, A. Lamperti, R. Mantovan, B. Ockert, E. E Fullerton, and D. Ravelosona, Low depinning fields in Ta-CoFeB-MgO ultrathin films with perpendicular magnetic anisotropy, Appl. Phys. Lett. 103, 182401 (2013), 10.1063/1.4826439 [4] M.Kisielewski, Z.Kurant, M.Tekielak, W.Dobrogowski, A.Maziewski, A.Wawro, L.T.Baczewski, Magnetooptical micromagnetometry of ultrathin Co wedges, Physica Status Solidi (A), vol.196, No.1, 129-132 (2003), 10.1002/pssa.200306370. [5] J. Ferré, V. Grolier, P. Meyer, S. Lemerle, A. Maziewski, E. Stefanowicz, S. V. Tarasenko, V. V. Tarasenko, M. Kisielewski, and D. Renard, Magnetization-reversal processes in an ultrathin Co/Au film, Phys Rev B 55, 15092–102 (1997), 10.1103/PhysRevB.55.15092. [6] A. K. Dhiman, T. Dohi, W. Dobrogowski, Z. Kurant, I. Sveklo, S. Fukami, H. Ohno, A. Maziewski, Magnetization processes and magnetic domain structures in Ta/CoFeB/MgO stacks, accepted in JMMM. [7] N. Vernier, J.-P. Adam, S. Eimer, G. Agnus, T. Devolder, T. Hauet, B. Ocker, F. Garcia, and D. Ravelosona, Measurement of magnetization using domain compressibility in CoFeB films with perpendicular anisotropy, Appl. Phys. Lett. 104, 122404 (2014), 10.1063/1.4869482

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/b5d1c9989052d51f4a6fef7bbdf051ff.jpg)
Fig. 1 The dependence of normalized magnetization m (derived from domains area) on total duration of magnetic field pulses t for the stack with d= 1.24 nm. The amplitudes of magnetic field pulses in series are 0.75, 0.85 and 0.95 mT.

![](https://assets.underline.io/uploads/markdown_image/1/image/7b3b405252fcac9a2760f9498043027a.jpg)
Fig. 2 The dependence of the normalized magnetization m (left Y-axis) and total length of NSDs, extracted from the field of view 1.4 mm2 (right Y-axis) on amplitude of magnetic field pulses m0H (pulse duration 500 ms) for the stacks with d=1.24nm.

 Magnetization Processes and Magnetic Domain Structures in Ta/CoFeB/MgO Stacks

Magnetic properties of an ultra-thin CoFeB film, particularly at the CoFeB/MgO interface, are of critical importance for the development of magnetic tunnel junction based magnetoresistive random access memory (MRAM) such as spin transfer torque MRAM (STT-MRAM) and voltage-controlled MRAM (VC-MRAM), because these MRAM cells have a CoFeB free layer with a thickness of around 1 nm in direct contact with a MgO tunnel barrier layer. In STT-MRAM cells, the CoFeB free layer is magnetized perpendicularly to the film plane by taking advantage of interfacial perpendicular magnetic anisotropy (PMA) induced at the interface with MgO [1,2]. Larger interfacial PMA is advantageous for shrinking the cell size while maintaining a sufficiently large thermal stability factor of the free layer magnetization. In VC-MRAM cells, the interfacial PMA is not only used to perpendicularly magnetize the free layer, but also actively modulated by the application of electric fields to trigger the so-called precessional toggle switching [3]. Therefore, thorough understanding of the physical mechanism of interfacial PMA is indispensable for material/process engineering of these MRAM technologies. The source of the interfacial PMA is considered to be preferential occupation of the dxy electron orbit that lies in-plane due to broken symmetry at the interface and spin-orbit interaction [4,5], which revives the orbital angular moment of the d-electrons that is almost completely quenched in the bulk. The magnitude of the orbital angular moment can be evaluated by the deviation of spectroscopic g factor from 2, i.e., g=2+L/S, where S and L are the spin and orbital angular moments, respectively. It has been observed that the magnitude of interfacial PMA is strongly dependent on the annealing temperature, suggesting that the crystalline state affects the electronic structure at the CoFeB/MgO interface, thus changing the interfacial PMA. According to these considerations, the spectroscopic g factor is expected to increase at the CoFeB/MgO interface with an annealing temperature dependence similar to that of the interfacial PMA. To examine this notion, we studied static and dynamic magnetic properties of CoFeB/MgO film stacks as a function of CoFeB thickness and annealing temperature. The samples are Ta(5) / CoFeB (t) / MgO(2) / Ta(5) / Ru(5) film stacks sputter deposited on a thermally oxidized Si substrate (from bottom to top, number in parentheses is the thickness of each layer in nm), where t is the nominal CoFeB thickness ranging from 1.1 nm up to either 1.5 or 5 nm for different annealing temperatures. These films were left as deposited or annealed between 150°C and 350°C for 1 hour in vacuum. We performed vibrating sample magnetometry (VSM) measurement for static characterization, and broadband ferromagnetic resonance (FMR) measurement for dynamic characterization, respectively. The CoFeB films have a magnetically dead layer at the Ta/CoFeB interface due to Ta diffusion with thicknesses depending on the annealing temperature, and the saturation magnetization Ms also varies on annealing temperature due to different degrees of crystallization. From the static M-H loop obtained by VSM, Ms as a function of t was first plotted, from which the effective thickness teff (t minus dead layer thickness) and effective Ms of the CoFeB layer were estimated for each annealing temperature. Next, the total PMA energy was calculated by integrating the static M-H loop. The interfacial PMA energy was estimated from the linear regression on the plot of total PMA energy as a function of teff for each annealing temperature. Figure 1 shows the interfacial PMA energy as a function of annealing temperature, which becomes maximum at 250°C. In the dynamic magnetic characterization, the broadband FMR measurement performed under out-of-plane bias magnetic field HB observed a Kittel mode FMR spectrum. The FMR frequency showed a liner relation on HB, from the slope of which the spectroscopic g factor was determined. The value of g factor showed a significant increase as teff decreased. Figure 2 shows the increase of g factor at the CoFeB/MgO interface from the bluk value as a function of annealing temperature, which becomes maximum at 200°C, and exhibits a very similar dependence on the annealing temperature to the interfacial PMA. These results support the notion that the increase of g factor and interfacial PMA have the same physical origin, i.e., revival of the orbital angular moment due to broken symmetry at the CoFeB/MgO interface. In the talk, we will present the detailed experimental and analytical procedures for studying the static and dynamic magnetic properties of the ultra-thin CoFeB film at MgO interface performed in this work, and possible physical origin to explain the observation results. This presentation was partly based on results obtained from a project, JPNP16007, commissioned by the New Energy and Industrial Technology Development Organization (NEDO), Japan, and JSPS KAKENHI 18K14118.

**References:**

[1] S. Yakata et al., J. Appl. Phys. 105, 07D131 (2009). [2] S. Ikeda et al., Nat. Mater. 9.9, 721 (2010). [3] Y. Shiota et al., Nat. Mater. 11.1, 39 (2012). [4] S. Miwa et al., J. Phys. D 52.6, 063001 (2018). [5] B. Dieny and M. Chshiev, Rev. Mod. Phys. 89.2, 025008 (2017).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/448829d743d9b248d164789bdfd98766.jpg)
Fig. 1, Interfacial PMA at CoFeB/MgO interface as a function of annealing temperature.

![](https://assets.underline.io/uploads/markdown_image/1/image/4b735ce897f2db9be15f9a3abeeb48b4.jpg)
Fig. 2, Increase of g factor at CoFeB/MgO interface as a function of annealing temperature.

 Effect of Annealing Temperature on Spectroscopic g Factor at CoFeB/MgO Interface

Bulk perpendicular magnetic anisotropy materials are proposed to be a promising candidate for next-generation ultra-high density and ultralow energy-consumption spintronic devices (1~5). In this presentation, we will report the experimental investigation on the structure, thermal stability, and magnetic properties of FePd thin films seeded by a Ru layer. An fcc-phase Ru layer induces the highly-ordered L10-phase FePd thin films with perpendicular magnetic anisotropy (Ku ~ 10.1 Merg/cm3). The thermal stability of FePd samples is then studied through the annealing process. It is found that a Ku ~ 6.8 Merg/cm3 can be obtained with the annealing temperature of 500 oC. In addition, the damping constant α, an important parameter for switching current density, is determined as a function of the testing temperature. We observe that α increases from 0.006 to 0.009 for the as-deposited FePd sample and from 0.006 to 0.012 for the FePd sample annealed at 400 oC as the testing temperature changes from 25 oC to 150 oC (6). These results suggest that Ru-seeded FePd provides great potential in scaling perpendicular magnetic tunnel junctions below 10 nm for applications in ultralow energy-consumption spintronic devices.

**References:**

 1. M. Weisheit, S. Fähler, A. Marty, Y. Souche, C. Poinsignon, and D. Givord, Science 315, 349-351 (2007). 2. S. Iihama, A. Sakuma, H. Naganuma, M. Oogane, T. Miyazaki, S. Mizukami, and Y. Ando, Appl. Phys. Lett. 105, 142403 (2014). 3. S. Iihama, A.Sakuma, H. Naganuma, M. Oogane, S. Mizukami, and Y. Ando, Phys. Rev. B 94, 174425 (2016). 4. D.-L. Zhang, K. B. Schliep, R. J. Wu, P. Quarterman, D. R. Hickey, Y. Lv, X. Chao, H. Li, J.-Y. Chen, Z. Zhao, M. Jamali, K. A. Mkhoyan, and J.-P. Wang, Appl. Phys. Lett. 112, 152401 (2018). 5. D.-L. Zhang, C. Sun, Y. Lv, K. B. Schliep, Z. Zhao, J.-Y. Chen, P. M. Voyles, and J.-P. Wang, Phys. Rev. Applied 9, 044028 (2018). 6. D.-L. Zhang, D. Huang, R. J. Wu, D. Lattery, J. Liu, X. Wang, D. B. Gopman, K. A. Mkhoyan, J.-P. Wang, and X. Wang, Appl. Phys. Lett. 117, 082405 (2020)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/759ad544c5fd3f1e39164b751e71e26b.jpg)
FIG. 1. (a) Atomic-resolution HAADF-STEM image of the as-deposited Ru-seeded FePd thin film. (b) The out-of-plane (θ-2θ scans) XRD patterns as a function of the annealing temperatures for Ru-seeded FePd thin films. The very clear (001) and (002) peaks of L10-phase are observed, indicating the high-ordered FePd thin films. (c), (d) The temperature-dependent effective anisotropic field (Hk,eff) and damping constant extracted from TR-MOKE measurements for both as-deposited and 400 oC-annealed samples.

 Low Gilbert damping and high thermal stability of Ru-seeded L10-phase FePd perpendicular magnetic thin films at elevated temperatures

Investigation of a creep of domain walls (DWs) in a combination of the in-plane and out-of-plane magnetic fields may give useful information about the Dzyaloshinskii-Moriya interaction in magnetic films. The first approach of the definition of the DMI field implied simple searching for the minimum in the dependence of the velocities of the DWs on the in-plane magnetic field, v(Hx) [1]. It was considered that when the in-plane magnetic field compensates the DMI field, the energy density of the DW is maximal and the velocity of the DW is minimal. However, papers were published in which the DMI fields measured by other methods like Brillouin light spectroscopy were very different from the magnetic fields of minima in the velocity curves v(Hx) [2]. To explain the shape of strongly asymmetric v(Hx) dependencies different extended models of the creep of DWs were inroduced [3, 4]. In this study, we investigate the DMI in the epitaxial Si(111)/Cu/Pd/Co(0.7 nm)/Pd system with the thickness of the bottom Pd layer varying from 0 to 3 nm. Since we measured non-zero DMI in the symmetric Pd/Co/Pd system in the previous paper [5], present study is aimed to understand the origin of the DMI in the symmetric Pd/Co/Pd system. The velocity curves v(Hx) were fitted by the extended dispersive stiffness model [4, 6] and the energy of DMI was derived from the fittings. We carried out the investigation of strains in the interfaces depending on the thickness of the Pd bottom layer, dPd, by a reflection high energy electron diffraction during the process of fabrication of the samples. We found a large asymmetry of the strains between the bottom Pd/Co and top Co/Pd interfaces in the samples. The bottom Pd layer is initially largely strained when it grows on the Si/Cu(2 nm) surface. Stress is gradually relaxed by the introduction of misfit dislocations in the Pd layers. Co layers grown on the top of Pd underlayers are strained. Strains are partially relaxed during growth of the Co layers. Therefore, strains in the bottom and in the top of Co layers depend on the thickness of the Pd bottom layer. Magnetic hysteresis loops of the samples were measured by a vibrating sample magnetometer and the energy of a perpendicular magnetic anisotropy was defined in the series of samples. We investigated the creep of domain walls by means of a Kerr-microscope eqipped with in-plane electromagnet and home-made perpendicular coil. Measured velocity curves of the samples were strongly asymmetric. They may be fitted well only by using a variable velocity prefactor in the creep law, v = v0exp(-ΔE/kT), v0 = v0*(1 + χ*cos(φeq)), where ΔE is an energy barrier, χ* is the weight of a chiral damping effect, φeq is an azimuthal angle of the magnetization in the DW. Strictly speaking the nature of the dependence of the velocity prefactor on the in-plane magnetic field is still under debate. In some studies, this effect is named chirality-induced or additional asymmetry effect. The dependence of the chiral damping weight on the Pd bottom layer thickness is shown in Fig. 1. The absolute value of the chiral damping weight increases up to 0.95 at the thickness of the bottom Pd layer of 1 nm and oscillates near this value with further increasing of the thickness of Pd. This behavior of the chiral damping weight is correlated with the dependence of the asymmetry of the strains between the bottom and top interfaces on the thickness of the Pd bottom layer. Therefore, we assume that the possible origin of the strong asymmetry of the velocity curves may be related with strains gradient in the Co layers. The energy of the effective DMI, Deff, as a function of the thickness of the bottom Pd layer is outlined in Fig. 2. The negative energy of effective DMI slightly increases when dPd = 0.2 nm. It may be explained by a structural inhomogeneity of the bottom interface and possible intermixing of Pd and Cu. With further increasing of dPd, the magnitude of |Deff| decreases. Deff changes a sign at the thickness of Pd layer of 0.8 nm and increases with an increase in the thickness of the Pd bottom layer. The overall dependence of Deff on the dPd may be explained by the dependence of strains in the interfaces on the thickness of Pd bottom layer. A kink in the Deff(dPd) dependence correlates well with the kink in the dependence of the elastic strains on the dPd. For the first sight the change in the sign of Deff may be easily explained by the growing influence of the contribution to the DMI from the bottom Pd/Co interface. However, we prepared asymmetric Pd(3 nm)/Co(0.7 nm)/Cu sample with the same strains distribution as in the Pd(3 nm)/Co(0.7 nm)/Pd sample but with near zero DMI. Therefore, the non-zero DMI in the symmetric Pd(3 nm)/Co/Pd(3 nm) sample is caused mainly by the contribution to the DMI from the top Co/Pd interface. We may conclude that unequal elastic strains in the bottom Pd/Co and top Co/Pd interfaces cause non-zero DMI in symmetric Pd/Co/Pd system. Authors acknowledge the financial support of the Russian Ministry of Science and Higher Education under the state task (0657-2020-0013), the Grant program of the Russian President (Grant No. MK-1384.2021.1.2) and Russian Foundation for Basic Research according to the research project No.19-02-00530.

**References:**

1. S.-G. Je, D.-H. Kim and S.-C. You, Phys. Rev. B., Vol. 88, p. 214401 (2013). 2. R. Soucaille, M. Belmeguenai, J. Torrejon, Phys. Rev. B., Vol. 94, p. 104431 (2016). 3. J. P. Pellegren, D. Lau and V. Sokalski, Phys. Rev. Lett., Vol. 119, p. 027203 (2017). 4. D.M.F. Hartmann, R. A. Duine, M.J. Meijer, Phys. Rev. B., Vol. 100, p. 094417 (2019). 5. A.V. Davydenko, A.G. Kozlov, A.G. Kolesnikov, Phys. Rev. B., Vol. 99, p. 014433 (2019). 6. D. Lau, J. P. Pellegren, H. T. Nembach, Phys. Rev. B., Vol. 98, p. 184410 (2018).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/616d83d450af5977728230e236ec3566.jpg)
Fig. 1. The Pd thickness dependence of the chiral damping weight in the Si(111)/Cu/Pd(0–3 nm)/Co(0.7 nm)/Pd(3 nm) samples.

![](https://assets.underline.io/uploads/markdown_image/1/image/5a613be85fa0118433cb6bb5ec4ab78d.jpg)
Fig. 2. The Pd thickness dependence of the energy of the effective DMI in the Pd/Co/Pd films.

 Creep of domain walls in epitaxial Pd/Co/Pd(111) trilayers

Permanent micro-magnets are widely used in many devices such as MEMS and diverse biological applications. However, their fabrication typically includes several complex processing steps. MnBi is an intermetallic ferromagnetic alloy that possesses good permanent magnet properties such as a high Curie temperature of 711 K, a high magnetocrystalline anisotropy of 1.6 MJ/m3 at room temperature, a relatively high (BH)max of around 20 MGOe and, remarkably, an unusual positive temperature coefficient of the coercivity [1,2], which makes MnBi an attractive rare earth-free alternative for high temperature applications. In this study, a one-step route has allowed for the fabrication of highly crystalline hexagonal MnBi micro-islands (lateral dimension of 30 µm and thickness of 0.2 – 2.1 µm), by RF-sputtering of a composite target onto glass substrates at different deposition temperatures (TD) [3]. A strong impact of TD on the magnetic properties has been proven. Specifically, magnetic anisotropy could be tuned from out-of-plane (TD &lt; 375 K) to in-plane (TD = 475 K), remarkably, without applying any external magnetic field during the deposition. High values of coercivity have been obtained: 13.1 kOe at 400 K along the out-of-plane direction for the sample grown at 375 K and 14.1 kOe at 400 K along the in-plane direction for the sample grown at 475 K. The sample topography and the magnetic domain structure have been studied, observing a strong perpendicular anisotropy in micro-islands with a crater-like shape as shown in the scheme of Figure 1. It has been observed that the periodicity of stripe domains obeys the D = 1.18 * sqrt (T) law [4] in good agreement with theoretical predictions (data points and red line in Figure 1). We have validated the dependency of D vs T in MnBi thin films previously reported in literature, which has been possible through the MFM study (Fig. 1) of the crater-like shape of our MnBi micro-islands [3]. Moreover, this study has allowed us to extend the dependency of D vs T for MnBi in a thickness range well beyond previously published results. Acknowledgements IMDEA acknowledges financial support from the EU M-ERA.NET Network and MINECO through "NEXMAG" (Project Success Case at H2020 M-era.Net, Ref. PCIN-2015-126) and "3D-MAGNETOH "(Ref. MAT2017-89960-R); and from CM through "NANOMAGCOST" (Ref. P2018/NMT-4321).

**References:**

[1] K. Patel et al., Nanoscale 10, 11701 (2018) [2] E. Céspedes et al., J. Alloys Compd. 729, 1156 (2017); M. Villanueva et al., AIP Adv. 9, 035325 (2019) [3] M. Villanueva et al., Submitted for publication (2021) [4] C. Kittel, Physical Review 70, 11-12 (1946)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/65afd49c6f719bdfe8463c50b2dc14a0.jpg)
Figure 1. Dependence of domain width (D/2) on the thickness (T) of the MnBi micro-islands from MFM images (inset) measured at remanence after sample demagnetization. The D = 1.18 * sqrt (T) law is obtained by fitting the data points (red line).

 High-coercive hexagonal MnBi micro-islands with tunable magnetic anisotropy and stripe magnetic domain patterns

The III-nitride family of wide bandgap semiconductors GaN, AlN, and their alloys are important for diverse applications ranging from solid-state lighting to RF and power electronics [1]. The long spin lifetimes in nitride semiconductor platform [2] makes it attractive for exploiting the spin degree of freedom of conducting electrons. The ferromagnet/semiconductor heterostructure is crucial for fundamental spin-related device building blocks such as spin injection, spin transport, spin detection [3] and spin to charge conversion [4]. Memory devices utilizing spin-orbit torques require smooth interfaces between ferromagnets and the heavy metal or topological insulator layers on top for efficient spin transmission [5]. Epitaxial growth of magnetic layers with smooth surfaces on GaN hosting desirable properties for spintronic applications will provide a path towards spintronic devices for energy-efficient memory applications, and its integration with GaN-based RF, photonic, and wide-bandgap CMOS platforms [6]. Mn4N, a metallic nitride ferrimagnet, is an attractive candidate for direct epitaxial integration with GaN and AlN for all-nitride ferromagnet/semiconductor heterostructures. MBE grown Mn4N thin films on cubic substrates such as MgO and SrTiO3 (STO) exhibit desirable properties for spintronic applications such as a high critical temperature (TN ∼ 740 K), large spin polarization (P ∼ 70 %), strong perpendicular magnetic anisotropy (Ku = 1.1×105 J/m3), low saturation magnetization (Ms = 7.1×104 A/m on STO), large domains (∼ millimeter size on STO) and high domain wall velocities (up to 900 m/s) [7, 8]. Through exploration of nucleation and growth conditions, we uncover plasma-assisted MBE growth conditions needed for significantly improved epitaxial growth of c-axis aligned Mn4N on GaN with smooth surface morphologies. A series of four samples were grown in which the growth temperature of the 80 nm Mn4N layer was varied from Ts = 150 C to Ts = 375 C, after the deposition of 100 nm homoepitaxial undoped GaN buffer layers at Ts = 670 C, as shown in Fig. 1. Instead of out-of-plane [111] orientation of Mn4N on GaN as found in reactive MBE [9], Mn4N layers grown using plasma-assisted MBE in this work are dominated by [001] orientation, and exhibit 12-fold in-plane symmetry in the diffraction pattern. Smooth Mn4N layers are obtained at low growth temperatures of Ts &lt;= 300 C, though deep pits with depth ~ 10 nm measured by atomic force microscopy are present in the film grown at Ts = 300 C. The magnetic properties of Mn4N grown on hexagonal GaN are comparable to those in earlier reports on cubic substrates such as MgO, and can be tailored by varying the growth temperature. For example, with the increase of growth temperature, the anomalous Hall resistance hysteresis loop (Fig. 2) not only becomes squarer but also exhibits an interesting sign-flip from n-type to p-type between Ts = 225 C and Ts = 300 C. Details about the structural and magnetic properties of Mn4N on GaN will be presented.

**References:**

[1]. D. Jena, et al., Japanese Journal of Applied Physics 58, SC0801 (2019). [2]. S. Krishnamurthy, et al., Applied physics letters 83, 1761 (2003). [3]. A. Song, et al., Applied Physics Express 13, 043006 (2020). [4]. W. Stefanowicz, et al., Physical Review B 89, 205201 (2014). [5]. Q. Shao, et al., Nature communications 9, 1 (2018). [6]. S. J. Bader, et al., IEEE Electron Device Letters 39, 1848 (2018). [7]. T. Gushi, et al., Nano Letters 19, 8716 (2019). [8]. T. Gushi, et al., Japanese Journal of Applied Physics 57, 120310 (2018). [9]. S. Dhar, et al., Applied Physics Letters 86, 112504 (2005).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/99cbce60474038d00770cf3b06e8bed4.jpg)
Fig. 1 (a) Schematic of the epitaxial structures in this study and (b) Crystal and magnetic structure of Mn4N.

![](https://assets.underline.io/uploads/markdown_image/1/image/48b511882754c881e1a60b3bbff1eba4.jpg)
Fig. 2 Anomalous Hall resistance hysteresis loops at 300 K for Mn4N layers grown on MgO at (a) 450 C and on GaN at (b) 150 C (c) 225 C (d) 300 C and (e) 375 C. Note the flip in the loop orientation from (c) to (d), indicating n-type to p-type conversion.

 Epitaxial Ferrimagnetic Mn4N Thin Films on GaN by Molecular Beam Epitaxy

The ability to modify the magnetic anisotropy and to control spin state of AFM materials is interesting from both a fundamental and an application point of view. By taking advantage of AFM/FM coupling, magnetic anisotropy of AFM sublayer can be tailored either by appropriate choice of FM thickness or in uniform thickness system by temperature induced processes [1]. Additionally, magnetic state of the AFM can be controlled by small external magnetic field. Here, we follow the magnetic properties of epitaxial NiO(111)/Fe(110) bilayers on W(110) substrate using XMCD and XMLD spectroscopy together with PEEM microscopy. We find that the magnetic spins in NiO are rotatable within the NiO(111) plane and both NiO magnetic anisotropy and orientation of AFM spins are determined by the magnetic properties of underlying ferromagnetic Fe layer. Specifically, magnetic state and anisotropy of antiferromagnetic NiO can be modified by tuning the thickness of Fe, changing the temperature or applying small external magnetic field. A special sample was prepared for X-PEEM measurements in order to image the magnetic domain pattern below and above the critical Fe thickness of in-plane Spin Reorientation Transition (SRT). The sample consisted two 50Å and 200Å thick Fe(110) regions, covered by homogeneous (40Å) layer of NiO(111). Fig.1 a shows XMCD-PEEM and XMLD-PEEM images of the border spot between the two Fe thicknesses. The photon helicity k is aligned parallel to the Fe[1-10] direction which is easy axis of thin Fe and perpendicular to the bulk-like Fe[001] easy axis of 200Å thick Fe(110) region. The zig-zag pattern indicates the 90° in-plane rotation of Fe magnetization across the Fe (50/200Å) boundary. XMLD-PEEM image acquired at the Ni L2 edge (Fig.1 a, middle panel, “XMLD Ni L2V POL”), with the linear polarization of incoming X-rays oriented in the Fe(110) || NiO(111) surface plane shows that magnetic domain structure of Fe in the vicinity of SRT is directly imprinted into the antiferromagnetic NiO overlayer. The transition between orthogonally oriented ferromagnetic Fe regions is reflected in the shape of the magnetic domain structure of AFM NiO. In XMLD-PEEM image acquired with horizontal polarization of incoming beam (bottom panel of Fig.1a, “Ni L2 H POL”) magnetic contrast totally vanishes because the X-ray polarization is along the surface normal. This confirms that the NiO spins are confined within NiO(111) plane and no out-of-plane magnetic components are present. In case of NiO, the magnitude of XMLD effect is usually defined by the ratio of the two characteristic peaks at L2 absorption edge of Ni (RL2). We define RL2 as the higher-energy peak intensity divided by the intensity of the lower one. RL2 ratios were determined for both normal (γ = 0○) and grazing (γ = 60○) incidence geometries and we define their difference as |ΔR L2| = |RL2(γ = 0○) − R L2 (γ = 60○)|. Fig.1 b shows polar plots ΔR L2(φ) of ΔRL2 as a function of the azimuthal in-plane angle φ between polarization and Fe[1-10] direction. Instead of three-fold in-plane symmetry observed for bulk NiO(111) in pioneering works [2], exchange coupling with neighboring Fe layer makes the in-plane magnetic anisotropy of NiO two-fold with its easy axis determined by the magnetic anisotropy of ferromagnet. Specifically, NiO and Fe spins couple collinearly at the interface and so the easy axis of NiO is parallel to Fe[1-10] or Fe[001] on 50 and 200Å thick regions, respectively. In Fig. 1 c we show element sensitive magnetic hysteresis loops measured by means of XMCD and XMLD for ferromagnetic Fe and antiferromagnetic NiO sublayers, respectively. Typical square hysteresis curve is observed for thin Fe region which is below SRT critical thickness, while hard axis loop with almost zero magnetization in remanent state is presented for 200Å thick Fe. In case of NiO, on 50Å thin Fe region, external magnetic field makes no changes in XMLD as it is insensitive to 180° reversal of AFM spins. Contrary, on thick Fe one can clearly see a change of RL2 as a function of small external magnetic field (~ 500 Oe). Recently we showed that in a uniform thickness NiO(111)/Fe(110) epitaxial bilayer system, at given temperature near 300 K, two magnetic states with orthogonal spin orientations can be stabilized in antiferromagnetic NiO [1]. Field-free, reversible switching between these two AFM states was demonstrated. Here we present results of systematic investigations of such field-free, temperature induced switching of AFM NiO spins. A sample prepared for that purpose contained several 300 µm wide Fe stripes with different thickness in the range of (92Å – 113Å). Fig. 2 shows temperature dependence of RL2 ratio for various thicknesses of Fe (only heating branch is shown for clarity). It is clear that depending on the Fe thickness a critical temperature of SRT in AFM NiO can be tuned. With increasing Fe thickness the critical temperature, at which NiO spins rotate in-plane by 90○ towards Fe[1-10] direction, continuously increases. This provides possibility to cover wide temperature window ~ (250 – 380 K) for field-free, reversible switching of AFM spins.

**References:**

References: [1] M. Slezak, P. Drozdz, W. Janus, H. Nayyef, A. Koziol-Rachwal, M. Szpytma, M. Zajac, T.O. Mentes, F. Genuzio, A. Locatelli and T. Slezak. Nanoscale 2020, 12, 18091. [2] K. Kurosawa, M. Miura, S. Saito, J. Phys. C: Solid St. Phys. 13 (1980) 1521-7

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/4219bb73c0734141604e4caa280a32fd.jpg)
Fig. 1 (a) XMCD-PEEM and XMLD-PEEM images at the border between thick (up) and thin (down) Fe (110) regions (200 and 50Å). (b) polar plot of XMLD for NiO (111) sublayer as a function of azimuthal in-plane angle (c) XMCD and XMLD magnetic hysteresis loops for Fe and NiO respectively.

![](https://assets.underline.io/uploads/markdown_image/1/image/7cf20ee7c47f47142b421c6a84a02f0c.jpg)
Fig. 2. Temperature dependence of XMLD (RL2) shows SRT in AFM NiO during heating process for various Fe(110) thicknesses.

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Tailoring the magnetic anisotropy and controlling the spin orientation in antiferromagnetic NiO(111) films on Fe(110).

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