Tuning interfacial Dzyaloshinskii-Moriya interactions in thin amorphous ferrimagnetic alloys

Skyrmions can be stabilized in magnetic systems with broken inversion symmetry and chiral interactions, such as Dzyaloshinskii-Moriya interactions (DMI). Further, compensation of magnetic moments in ferrimagnetic materials can significantly reduce magnetic dipolar interactions, which tend to favor large skyrmions. Tuning DMI is essential to control skyrmion properties, with symmetry breaking at interfaces offering the greatest flexibility. However, in contrast to the ferromagnet case, few studies have investigated interfacial DMI in ferrimagnets. Here we present a systematic study of DMI in ferrimagnetic CoGd films by Brillouin light scattering. We demonstrate the ability to control DMI by the CoGd cap layer composition, the stack symmetry and the ferrimagnetic layer thickness. The DMI thickness dependence confirms its interfacial nature. In addition, magnetic force microscopy reveals the ability to tune DMI in a range that stabilizes sub-100 nm skyrmions at room temperature in zero field. Our work opens new paths for controlling interfacial DMI in ferrimagnets to nucleate and manipulate skyrmions.


Results
tuning the DMi with the capping layer composition. The ferrimagnetic CoGd thin films were grown by RF magnetron co-sputtering on oxidized silicon wafers in the following sequence: W(3)/Pt(3)/Co 78 Gd 22 (t)/ Pt 1-x W x (3)/Pt(3) (thicknesses in nanometers) [Methods]. The W/Pt seed layer provides good adhesion to the substrate and texture to ensure perpendicular magnetic anisotropy (PMA). The top Pt layer prevents sample oxidation. The DMI of a 5-nm thick CoGd film was studied as a function of the W composition (x) of the cap layer Pt 1-x W x . The magnetic properties of the films were measured by vibrating sample magnetometry (VSM) and are summarized in Table 1 (Methods). Figure 1(a) and 1(b) show an out-of-plane field room-temperature Spin wave spectroscopy using BLS was performed to measure the DMI in the CoGd films (Methods). The DMI leads to an asymmetric frequency dispersion of the counterpropagating spin waves 30 . The DMI energy D (mJ m −2 ) is proportional to the frequency shift (∆f DMI ) and given by: where g is the spectroscopic splitting factor (we take g = 2), μ B the Bohr magneton and h Planck's constant and k = 16.7 µm −1 is the spin wave vector. Notably, the DMI energy given by BLS measurements is an effective DMI averaged over the film thickness, i.e., a sum of the bottom and top interfacial contributions. An example of BLS spectra is displayed in Fig. 1(c) for Pt/CoGd(5 nm)/W. We fitted the spectra for positive and negative field polarity. The frequency shift was determined for the Stokes and the anti-Stokes peaks separately and then averaged. The diameter of a skyrmion results from the competition between different energies such as the Heisenberg exchange energy, the magnetic anisotropy energy and the DMI strength. Ultrasmall skyrmions can be nucleated at room temperature only in a narrow range of D [32][33][34] . For DMI strength larger than a scale set by the magnetic anisotropy, the formation of stripe domains become energetically favorable 33,34 . Conversely, a weak DMI cannot stabilize a skyrmion. Theoretical work has predicted that ferrimagnetic materials are better candidates than ferromagnets to host ultrasmall and ultrafast skyrmions due to their low saturation magnetization, which causes only small stray fields 22 . Indeed, in ferrimagnetic materials, the interfacial DMI can dominate over the dipolar interactions and enable the formation of ultrasmall DMI skyrmions, which is difficult to achieve in ferromagnets. Hence, our goal was to provide a new method for controlling the interfacial DMI in thin ferrimagnetic CoGd films, which could allow one to precisely tune the DMI in a range that would enable skyrmion nucleation.
Changing the nature of the CoGd interfaces can be used to engineer the DMI strength. Therefore, the idea is to leave the Pt underlayer at the bottom interface of the CoGd film unchanged and insert a Pt 1-x W x alloy at the top interface. Thus, by changing the composition of the Pt 1-x W x alloy, the structural symmetry of the film can be gradually broken to induce DMI. Pt is chosen for its strong spin-orbit coupling that gives a large interfacial DMI on Co spins 35,36 . Theoretical calculations based on Hund's first rule have shown that, on the contrary, a weaker DMI arises from interactions between W and Co 35 and with the same chirality as Pt and Co. In addition, W, due to its giant spin-Hall angle 37,38 , would serve as a spin current source to enable skyrmion motion induced by spin-orbit torque (SOT). Figure 2(a) shows the DMI energy as a function of the W composition (x) in Pt/CoGd(5 nm)/Pt 1-x W x measured by BLS. A maximum DMI of 0.23 ± 0.02 mJ m −2 is obtained for the asymmetric stack (x = 1). In comparison, a bulk DMI of up to 0.10 mJ m −2 was reported in ferrimagnetic GdFeCo films 26 . Conversely, the DMI is small for the symmetric film (x = 0). In the Pt/CoGd/Pt film, the top and bottom interfaces induce an interfacial DMI of similar amplitude but opposite sign (as the DMI is a chiral interaction), thus, resulting in a near vanishing effective DMI. In Pt/CoGd/W, since W gives rise to a weaker interfacial DMI, the contributions of the two interfaces are not compensated, leading to a larger effective DMI. As seen in Fig. 2(a), as little as 20% of W introduced in the top layer is enough to significantly break the symmetry and induce a DMI of 0.20 mJ m −2 . Yet, for x > 0.2, the DMI is less sensitive to the W content in the top interface. Nevertheless, for 0.1 < x < 1, the DMI increases by 50% to reach the maximum value of 0.23 ± 0.02 mJ m −2 . This would indicate that for 0.1 < x < 0.9, the W www.nature.com/scientificreports www.nature.com/scientificreports/ rather greatly reduces the DMI between the Co spins and the Pt. Indeed, if the W were actively contributing to the interfacial DMI, a stronger dependence of the DMI energy with the alloy composition would have been expected. Additionally, the quality of the interfaces, which has a great impact on the DMI 39,40 , was assessed by cross-sectional TEM. Figure 2(b) shows a cross-section of the asymmetric Pt/CoGd(5 nm)/W film, while a closer view of the top and bottom interfaces of the CoGd layer is displayed in Fig. 2(c). The CoGd alloy and the W layers are amorphous and the Pt is polycrystalline. Figure 2(c) shows that the Pt/CoGd and the CoGd/W interfaces are smooth.
thickness-dependence of the DMi. It is necessary to study the dependence of the DMI on CoGd thickness to establish its nature, i.e. to know whether the DMI is arising from interfacial effects. In our Pt/Co 78 Gd 22 (t)/ Pt 1-x W x films, the W composition (x) was fixed either to 0 or 1 to investigate the DMI in a symmetric (x = 0) and asymmetric (x = 1) stack as a function of the magnetic thickness t. t was increased from 5 nm to 15 nm. The magnetic properties were systematically measured by VSM as a function of thickness. The results are presented in Fig. 3(a) and the DMI is plotted versus the inverse magnetic thickness. In the asymmetric Pt/CoGd/W stack, the DMI is inversely proportional to the magnetic thickness and reaches a minimum of 0.09 ± 0.01 mJ m −2 for 1/t = 0.067 nm −1 (t = 15 nm). The DMI linearly increases with the inverse thickness. This confirms that the strength of the DMI at the interface remains unchanged and underlines its interfacial nature. The slope of the linear fit corresponds to the surface DMI constant D S ≃ 1 pJ m −1 , which is in the same order of magnitude than the reported surface DMI in ferromagnets 41,42 . In Fig. 3(b), the saturation magnetization times the magnetic thickness is plotted versus the CoGd thickness. It has a linear dependence on thickness with an x-axis intercept near zero thickness, which indicates that there is no measurable dead layer in the CoGd film. Notably, in Fig. 3(a), the intercept of the linear fit is non-zero for 1/t = 0 (i.e. an infinitely thick film). This indicates that there is a residual DMI of 0.025 mJ m −2 , which may result from a change of the magnetization compensation temperature throughout the thickness as evidenced in another rare-earth transition-metal alloy 43 that could induce inversion symmetry breaking. Yet, as the thickness decreases, the interfacial effects become more important and the DMI increases as seen in Fig. 3(a). Thus, the interfacial DMI dominates in the entire thickness range we have studied.
On the other hand, for the symmetric Pt/CoGd/Pt film, the DMI remains constant over the studied thickness range at about 0.10 ± 0.01 mJ m −2 as seen in Fig. 3(a) (red data points). This behavior is surprising as the interfacial DMI is expected to be almost zero in symmetric layer structures. This result shows there is a difference in the nature of the top and bottom CoGd interfaces. In order to verify the latter, we performed TEM imaging in the Pt/CoGd(15 nm)/Pt film. The full stack is shown in Fig. 4(a) and a closer view of the top and bottom interfaces in Figs. 4(b) and 4(c), respectively. In Fig. 4(b), a thin layer of intermediate gray contrast (indicated by the white arrows) can be seen at the top CoGd interface and not in the bottom interface. It appears that the Pt from the capping layer has diffused into the amorphous CoGd film. As a result, the bottom and top interfaces have different roughness and intermixing. Hence, the DMI contributions of the top and bottom interfaces are not equal. Thus, due to the chirality of the interaction, they do not cancel out, leading to an increase of the net DMI. Intermixing and roughness effects appear to be also dominant in thicker films as the DMI remains non-zero for larger thicknesses in Pt/CoGd/Pt as seen in Fig. 3(a). This is consistent with previous reports in the literature where the effect of interface roughness and intermixing on the inversion symmetry breaking has been extensively studied. It was reported that the DMI and the domain wall velocity increased with the difference of roughness and intermixing between the top and bottom interface of a magnetic layer 39,40,44,45 . Finally, as the skyrmion size depends on the magnetic film thickness 22,34 , it is thus important to understand how the interfacial DMI scales with the thickness. evidence of magnetic skyrmions by MfM. Finally, we aimed to verify whether these thin ferrimagnetic alloy films would indeed host skyrmions. We focused on the asymmetric Pt/CoGd/W stacks as they are www.nature.com/scientificreports www.nature.com/scientificreports/ more promising for skyrmion motion via spin-orbit torque because of the giant spin-Hall angle of W 37,38 . In fact, in Pt/CoGd/Pt, the spin-orbit torques from the top and bottom interfaces would tend to cancel each other out. The Pt/CoGd/W films were subject to AC in-plane magnetic field demagnetization and imaged by atomic and magnetic force microscopy (AFM and MFM) at room temperature in zero field.  www.nature.com/scientificreports www.nature.com/scientificreports/ is indicated by dark areas in the MFM images. By comparing the AFM and MFM images, it is clear that this contrast comes from magnetic textures and is not due to topography. Several skyrmion-like textures can be seen in Fig. 5(b). Figure 5(d) corresponds to a smaller MFM scan performed around of one of them marked by a square box in Fig. 5(b). This skyrmion-like texture is on the order of 100 nm. 50 nm skyrmions were observed in Pt/ CoGd(8 nm)/W (see supplemental materials). Arguably, considering the size of these textures, the DMI energy values [see Fig. 3(a)], and the fact that the CoGd films are weakly magnetized (M S ~ 140 -150 kA m −1 , see supplemental materials), it is unlikely that these textures are magnetic bubbles 4,7 stabilized by dipolar interactions. Thus, MFM images would rather indicate the presence of skyrmions. However, accurate estimation of the skyrmion size is difficult. Indeed, the MFM tip is sensitive to the dipolar field emerging from the magnetic texture which is spatially spread out. Furthermore, smaller magnetic features may be present in Fig. 5(b), yet they cannot be clearly distinguished due to the background noise and small magnetic contrast.

Discussion
To summarize, we have demonstrated that by capping the ferrimagnetic CoGd layer with a PtW alloy we could tune the DMI energy over a large range, from almost no DMI to an interfacial DMI energy of 0.23 mJ m −2 . The DMI thickness dependence reveals the interfacial nature of the DMI in CoGd thin films. Thus, the DMI strength can be controlled by the interfaces in the thickness range we studied, which is also the range relevant for skyrmion nucleation. Hence, interfacial DMI can not only be tuned by changing the cap layer composition but also by changing the thickness of the CoGd layer. Together they provide a wide range of tunability. However, the experimental results point out that changing the cap layer composition provides a lesser tunability than varying the thickness. Moreover, the DMI was found to be non-zero in thicker symmetric structures emphasizing the role of interface roughness and intermixing. Lastly, we showed evidence that films can have a DMI in a range that allows sub-100 nm skyrmion nucleation at room temperature in zero field. Our experimental results provide insight into the key parameters that control the DMI in ferrimagnetic films toward achieving ultrasmall and ultrafast skyrmion motion for spintronic applications.

Methods
Thin film deposition. The thin films were prepared by RF magnetron sputtering and deposited onto Si-SiO 2 substrates at room temperature with a base pressure of 2.7×10 −5 Pa. The Ar deposition pressures of W, Pt, CoGd, and Pt 1-x W x were 0.93 Pa, 0.1 Pa, 0.16 Pa, and 0.16 Pa, respectively. CoGd films were obtained by co-sputtering from the Co and Gd targets. The powers of the Co and Gd sources were tuned to obtain CoGd films with approximately 78 at. % of Co. The Pt 1-x W x alloy layer was also deposited by co-sputtering from the Pt and W targets. The alloy composition was varied by changing the deposition rate of the Pt and W targets. The deposition rates were calibrated using x-ray reflectometry.
Magnetometry. The magnetic properties of the samples were measured by vibrating sample magnetometry. Magnetic hysteresis loops were measured by varying the temperature from 100 K to 300 K with steps of 25 K in order to extract the temperature dependence of the saturation magnetization and the coercive field. Magnetometry was systematically performed prior to BLS experiments.
Brillouin light scattering. Spin wave spectroscopy using BLS is sensitive to interfacial effects and can be used to measure the DMI strength. The spin waves (SWs) inelastically scatter the monochromatic laser beam that is focused onto the sample surface. The frequency of the scattered photons is shifted by the SWs frequency. The SWs frequency is determined by analyzing the backscattered light with a (3 + 3)-pass tandem Fabry-Pérot interferometer. The counterpropagating Damon-Eshbach SWs have a non-reciprocal frequency dispersion characterized by a frequency shift (noted ∆f DMI ). The frequency shift is considered here in absolute value. An in-plane bias magnetic field sufficient to saturate the sample was applied to allow the SW to propagate in-plane (Damon-Eshbach geometry). For a λ = 532 nm laser beam with an incidence of θ i = π/4, the SW vector, k defined as k = 4πsin(θ i )/λ was set to 16.7 μm −1 . Damon-Eshbach SWs are almost plane waves extending throughout the thickness of the magnetic layer with their maximum amplitude reached at the top and bottom interface. For relatively thick layers, the photons may not reach the bottom interface due to the limited penetration depth into the sample structure. Nevertheless, the photons still interact with the SWs extending in the bulk of the layer and the energy transfer is the same and related to the interfacial DMI strength.