# Effect of oxide barriers and graphene intercalation on the magnetic properties of graphene/oxide/ferromagnet heterostructures.

I. Berrai¹, P. A. Dainone², M. Belmeguenai¹, Y. Roussigné¹, S. Tria³, N. Challab⁴, W. Alimi¹, Y. Lu², S. Farhat¹, S. M. Chérif¹*

¹Université Sorbonne Paris Nord, Laboratoire des Sciences des Procédés et des Matériaux, CNRS, LSPM-LPR 3407, 99, Avenue J.-B. Clément, 93430 Villetaneuse, France

²Institut Jean Lamour, Université de Lorraine, CNRS UMR 7198, Campus ARTEM, 2 Allée André Guinier, BP 50840, 54011 Nancy, France

³Université Sorbonne Paris Nord, Department of Physics, UFR SMBH, Campus de Bobigny, 1, rue de Chablis, 93017 Bobigny, France

⁴Université Paris Cité, Laboratoire Interfaces Traitements Organisation et DYnamique des Systèmes – CNRS, ITODYS-UMR 7086, 15, rue Jean de Baïf, 75013 Paris, France

## Abstract

We report on the effect of tunnel barrier intercalation between a monolayer Gr (MLGr) and a ferromagnet on structural, static and dynamic properties of SiO₂/MLGr/Oxide Barrier/Co₄₀Fe₄₀B₂₀/Ta heterostructures. Gr-based multilayers with MgO, elaborated by Molecular Beam Epitaxy, and Al₂O₃, elaborated by Atomic Layer Deposition, as tunnel barriers were investigated. Raman spectroscopy and vibrating sample magnetometry measurements were employed to evaluate the quality of Gr and the static magnetic properties of the material stacks, respectively. Additionally, Brillouin light scattering and micro-strip ferromagnetic resonance measurements were carried out to investigate the dynamic magnetic parameters, including the perpendicular magnetic anisotropy, the magnetic damping, and the interfacial Dzyloshinskii-Moriya interaction, whose control is essential for spintronic applications. Our study was completed by comparing the results with those obtained from control samples (without MLGr) and allowed discriminating the role played by the MLGr and pointing out the influence of oxide barriers on the magnetic properties of Gr/tunnel barrier/FM heterostructures.

**PACS:** 68.65.Pq, 75.30.Ds, 75.30.Gw, 75.70.Cn, 75.78.-n,

**Key words:** Graphene, Oxide barriers, CoFeB films, Perpendicular magnetic anisotropy, damping, Ferromagnetic resonance, Brillouin light scattering, spin waves and magnetization dynamics

* cherif@univ-paris13.fr
INTRODUCTION

The field of spintronic has undergone a remarkable evolution with the emergence of two-
dimensional (2D) materials [1]. Remarkably, among 2D materials, graphene (Gr) stands out
as exceptionally well suited for spintronic circuits due to its demonstrated ability for efficient
spin transport [2] and spin filtering [3]. For example, weak spin signals have been reported in
nanoscale lateral spin valves made of both metallic and semiconductor materials due to the
limited spin diffusion length [4]. Gr, on the other hand, offers the advantage of long spin
diffusion length ~100µm [5]. This property is essential for maintaining spin coherence to
develop spintronic devices such as spin valves [5,6], spin transistor [7]. These studies have
demonstrated that Gr can be successfully integrated into spin valve structures to enhance
electron spin manipulation and transport, paving the way for more efficient and promising
spin logic devices [8]. An interesting avenue of exploration in the field of spintronics is the
effect of integrating Gr into magnetic tunnel junctions (MTJs), fundamental building blocks
of spintronic devices consisting of two ferromagnetic (FM) layers separated by a thin
insulating layer, which play a crucial role in modern data storage and processing technologies,
mainly non-volatile magnetic random access memories (MRAMs). However, MTJs featuring
in-plane magnetic anisotropy exhibit limited thermal stability, high power consumption, and
manufacturing challenges. In response to these issues, research focus has shifted to MTJs
based on perpendicular magnetic anisotropy (PMA), a critical parameter for various magnetic
storage and spintronic devices, which is the preferred approach [9]. By coupling Gr with
ferromagnetic materials, a significant improvement in PMA has been demonstrated [10–13].
This effect is attributed to the strong coupling between the carbon atoms in Gr and the
ferromagnetic (FM) material atoms, leading to a modulation of the magnetic anisotropy
energy, mainly through the hybridization of the FM 3d orbitals and the Gr π bands [10–13].
Furthermore, it is noteworthy that achieving efficient spin signals requires careful control of
the interface between the 2D material and the ferromagnetic layer. In the case of Gr, in
addition to its intrinsic wavy surface, direct contact with a metallic layer causes defects or
even amorphization of the Gr layer, which reduces the effectiveness of the devices [14].
Berger et al. showed that direct contact between Co and Gr results in increased magnetic
disorder within the Co layers, potentially arising from alterations in the ferromagnet growth
process and an increased presence of structural defects [15]. This magnetic disorder is thought
to contribute to the observed broadening of the measured ferromagnetic resonance (FMR) line
width and consequently leading to an increase in the magnetic damping. Recently, changes of
PMA and damping in Gr/CoFeB-based heterostructures have been attributed to the presence
of the Gr ripples [16]. Introducing a high-quality tunnel barrier between Gr and the
ferromagnet not only resolves the impedance mismatch issue but also minimizes defects in the
ferromagnetic layer, serving as a critical factor in attaining both high magnetoresistance
(TMR) tunneling and efficient magnetization switching within MTJs [17]. In this context,
Al₂O₃ and MgO have become widely adopted as standard tunnel barriers, consistently
demonstrating their effectiveness in producing efficient tunneling devices [17,18]. The oxide
barrier and interface quality are influenced by the selection of materials and the deposition
methods. In the literature, molecular beam epitaxy (MBE) is classified as a physical vapor
deposition (PVD) method, while atomic layer deposition (ALD) is classified as a chemical
vapor deposition (CVD) one [19,20]. While these two methods are milder than magnetron
sputtering, which is also frequently employed for tunnel barrier deposition, it is noteworthy
that defects, such as oxygen vacancies and pinholes, are common in Al₂O₃ and MgO tunnel
barriers and can be challenging to avoid. This is particularly true when oxides are deposited
by PVD process. Furthermore, this process may introduce defects into 2D material on which
the barrier is deposited [17,21–23]. Although it is already reported for the comparison of
impact of sputtered Al₂O₃ and MgO tunnel barrier on the structural property of graphene [14],
the influence of barrier on the magnetic properties of Gr/tunnel barrier/FM heterostructure is
still not clear. Among these properties, perpendicular magnetic anisotropy (PMA), damping
and interface Dzyaloshinskii-Moriya interaction (iDMI) are of great interest for the next-
generation spintronic devices. For instance, in MTJs, large PMA can be used to obtain good
memory retention and a low write current can be achieved with a relatively weak
damping [24]. DMI is naturally present in structures with broken inversion symmetry. It
stabilizes Néel-type domain walls (DWs) and allows for nucleation and propagation of
topological textures such as skyrmions [25] by combining with PMA in thin films, which is
expected to be excellent candidate for the next generation of high density information storage
devices. It has also be shown that skyrmion-mediated switching of ferromagnets in presence
of thermal perturbation can be feasible down to lateral dimensions ~20 nm in MTJs thanks to
large values of voltage controlled magnetic anisotropy (VCMA) and DMI [26]. On the other
hand, it was reported that the DMI has a detrimental effect on spin-transfer torque (STT)
switching in conventional heavy metal/ferromagnet single-interface STT [27,28], but this
detrimental effect of DMI could be almost eliminated in double-interface STT-MTJs [29].

In this work, we report on the effect of tunnel barrier intercalation between a monolayer Gr (MLGr) and a ferromagnet on structural, static and dynamic properties of SiO₂/ MLGr/ Oxide Barrier (2 nm)/Co₄₀Fe₄₀B₂₀ (2 nm)/Ta (4 nm) heterostructures. Two Gr-based multilayers with MgO (by MBE) and Al₂O₃ (by ALD) as tunnel barriers were investigated. Raman spectroscopy and vibrating sample magnetometry (VSM) measurements were employed to evaluate the quality of Gr and the static magnetic properties of the material stacks, respectively. Additionally, Brillouin light scattering (BLS) and micro-strip ferromagnetic resonance (MS-FMR) measurements were carried out to investigate the dynamic magnetic parameters, including the perpendicular magnetic anisotropy (PMA), the magnetic damping, and the interfacial Dzyloshinskii-Moriya interaction (iDMI), whose control is essential for
applications. Finally, our study was completed by comparing the results with those obtained
from control samples (without MLGr), with the aim of discriminating the role played by the
MLGr and pointing out the influence of barrier on the magnetic properties of Gr/tunnel
barrier/FM heterostructure.

# I. EXPERIMENTAL RESULTS

## A. Sample Preparation, Structural and Static Magnetic Properties

A series of samples, Si/SiO₂/MLGr/Oxide Barrier (OxB: Al₂O₃ or MgO 2 nm)/Co₄₀Fe₄₀B₂₀ (2 nm)/Ta (4 nm) and Si/SiO₂/OxB (2 nm)/Co₄₀Fe₄₀B₂₀ (2 nm)/Ta (4 nm) (Fig.1(a)) were elaborated. The MLGr was synthesized in a homemade inductive heating CVD reactor with copper serving as the catalyst. Details have been presented elsewhere [30,31]. Once synthesized, the Gr is transferred to the Si/SiO₂ substrate by wet transfer method. The SiO₂/MLGr (and Si/SiO₂) samples were then annealed under ultra-high vacuum at 300°C for 1 hour to remove surface contamination, for the OxB deposition. The 2 nm of MgO barrier was deposited by MBE technique at 300°C at a speed of 0.02 nm/s, while the alumina was deposited by water based-ALD at the same temperature. Finally, CFB and Ta were deposited by magnetron sputtering with a power of 50 W and 100 W, respectively.

Raman spectroscopy was used to characterize the quality of graphene. All Raman
spectra were recorded at room temperature under laser excitation λexc = 473 nm. Figure.1(b)
shows the Raman spectrum of as-grown Gr, MLGr/Al₂O₃/CFB/Ta and MLGr/MgO/CFB/Ta
samples. One can observe the characteristics bands of graphene, mainly the primary in-plane
vibrational mode G-band at (~ 1580 cm⁻¹) which arises from the stretching of the C-C bond in
graphitic materials, 2D-band at (~ 2705 cm⁻¹), a second-order overtone of another in-plane
vibration and is highly sensitive to the number of graphene layers, and disorder-induced D-
band at (~ 1356 cm⁻¹) [32].
<figure><img src="image_1.png" /><figcaption>**Fig. 1.** (a) Scheme of the MLGr/OxB/CFB/Ta and OxB/CFB/Ta samples. (b) Raman spectra of the Gr-based samples. (c) Hysteresis loops of the Gr-based samples and controls.</figcaption></figure>

The number of layers can be established through the ratio of peak intensities I<sub>2D</sub>/I<sub>G</sub>, along with the quality of the 2D peak. In the depicted pristine graphene spectrum, the I<sub>2D</sub>/I<sub>G</sub> ratio is approximately 3.3. The 2D peak exhibits a Lorentzian shape with a singular component, and the full width at half maximum (FWHM) is approximately 31 cm<sup>-1</sup>. These specific characteristics serve as distinctive indicators of high-quality monolayer graphene [33]. It is noteworthy that the deposition of MgO on MLGr induces additional defects, resulting in an increase in the D peak intensity, as seen in the Raman spectrum of MLGr/MgO/CFB/Ta sample. In the case of Al<sub>2</sub>O<sub>3</sub>, this effect is less pronounced. It's important to emphasize that achieving ultrathin dielectric growth on sp<sup>2</sup>-bonded graphene layers presents significant challenges. Both preserving the original sp<sup>2</sup> structure of the graphene and ensuring uniform wetting of the dielectric require meticulous engineering efforts [23,34].
Figure.1(c) shows the normalized VSM hysteresis loops recorded at room temperature of all the samples. In the absence of MLGr, the samples demonstrate practically identical coercivity, indicating that the type of OxB has no discernible effect on the coercive $H_c$ and saturation $H_s$ fields of CoFeB. In the presence of MLGr, while $H_c$ appears to be little affected, the saturation occurs at higher $H_s$ fields compared to the controls and the squarness, ratio of remnant magnetization $M_r$ on saturation magnetization $M_s$, is clearly reduced.

## B. Dynamic Magnetic Properties

To investigate the PMA in presence of oxide barriers, while also highlighting the influence of graphene, we conducted BLS spectroscopy measurements using the Damon-Eshbach configuration [35]. This configuration enables the probing of spin waves (SWs) propagating along the in-plane direction, perpendicular to the applied in-plane magnetic field $H$. The measurements were performed with a low wave vector ($k_{sw} = 8.07$ µm$^{-1}$ i.e. $\theta= 20^\circ$) to prevent potential contributions caused by interfacial Dzyaloshinskii-Moriya interaction (iDMI). The results were compared with those obtained from in-plane FMR measurements. The Stokes $F_s$ (negative frequency shift) and anti-Stokes $F_{as}$ (positive frequency shift) SWs frequencies were determined through Lorentzian fits applied to the BLS spectra shown in Fig.2(a), while for the analysis of the in-plane FMR spectra shown in Fig.2(b), we used the same type of fit as for the out-of-plane FMR spectra (supplementary material [36]). It should be noted that the in-plane FMR spectra for MLGr/MgO/CoFeB structure present two modes at a given scanning frequency, whereas only one mode is observed for MLGr/Al₂O₃/CoFeB structure as shown in Fig.2(b). By analysing the BLS frequency variation with the in-plane applied field, we estimated the effective magnetizations $\mu_0 M_{eff} = \mu_0 (M_s - H_a)$ for each sample and its control, where $M_s$ is equal to (1216±50) kA/m, in agreement with Zhou *et*
al. [37], and $H_a$ is the PMA field which can be expressed as $H_a = \frac{2K_{eff}}{\mu_0 M_s}$ where $K_{eff}$ is the PMA constant, using relation (1).

$$F = \left(\frac{F_S+F_{as}}{2}\right) = \left(\frac{\mu_0\gamma}{2\pi}\right) \sqrt{[H+Jk_{sw}^2+P(k_{sw}t)M_s][H+Jk_{sw}^2-P(k_{sw}t)M_s+M_{eff}]} \quad (1)$$

Here, H represents the in-plane applied field, $\frac{\gamma}{2\pi}$ the gyromagnetic ratio about 30.3 GHz/T (29.7 GHz/T) for MLGr-Al₂O₃ (MLGr-MgO) based-system, respectively (supplementary material [36]), $\mu_0$ the vacuum permeability, $J = \frac{2A_{ex}}{M_s}$ the SW stiffness with $A_{ex}=10^{-11}$ J/m [38] the micromagnetic exchange constant of CFB, and $P(k_{sw}t) = 1 - \frac{1-\exp(-|k_{sw}|t)}{|k_{sw}|t}$ is a coefficient describing the dipolar interactions.

The results obtained from the in-plane applied field-dependent BLS frequencies depicted in Fig.2(c) demonstrate that in the absence of MLGr, both systems exhibit similar behavior and their effective magnetizations are nearly identical as shown in Table.1. Our data thus clearly suggest that there is no significant influence of the nature of the OxB on PMA in this case. The FMR measurements (Fig.2(d)) confirmed the BLS results as presented in Table.1. In the presence of MLGr, our results concluded to a reduction of the effective magnetization and a reinforcement of the PMA constant.
<figure><img src="image_2.png" /><figcaption>**Fig. 2.** (a) BLS Spectra of samples and controls at $\mu_0H = 0.3$ T and $k_{sw} = 8.07 \ \mu m^{-1}$. (b) FMR spectra of MLGr/Al₂O₃/CFB/Ta and MLGr/MgO/CFB/Ta samples at different frequencies. (c) In-plane applied field-dependent BLS frequencies. (d) FMR frequencies for samples and controls.</figcaption></figure>

<table><caption>**Table 1:** Effective magnetization, PMA constant values and damping coefficient as deduced from BLS and FMR results.</caption><thead><tr><th></th><th></th><th>SiO<sub>2</sub>/Al<sub>2</sub>O<sub>3</sub>/CFB</th><th>SiO<sub>2</sub>/MgO/CFB</th><th>MLGr/Al<sub>2</sub>O<sub>3</sub>/CFB</th><th>MLGr/MgO/CFB</th></tr></thead><tbody><tr><td rowspan="2">μ<sub>0</sub>M<sub>eff</sub>(T)</td><td>BLS</td><td>0.91 ± 0.02</td><td>0.97 ± 0.01</td><td>0.67 ± 0.01</td><td>0.76 ± 0.01</td></tr><tr><td>FMR</td><td>0.89 ± 0.01</td><td>0.97 ± 0.01</td><td>0.63 ± 0.01</td><td>0.83 ± 0.01<br>0.68 ± 0.02</td></tr><tr><td>K<sub>eff</sub></td><td>BLS</td><td>3.80 ± 0.65</td><td>3.4 ± 0.6</td><td>5.2 ± 0.6</td><td>4.7 ± 0.6</td></tr></tbody></table>
<table><tr><td>(x10<sup>5</sup> J/m<sup>3</sup>)</td><td></td><td></td><td></td><td></td></tr><tr><td></td><td>FMR</td><td>3.9 ±0.6</td><td>3.7 ±0.6</td><td>5.4 ±0.6</td><td>4.24 ±0.60</td></tr><tr><td></td><td></td><td></td><td></td><td></td><td>5.15 ±0.70</td></tr><tr><td>α (x 10<sup>-3</sup>)</td><td>BLS</td><td>7.6 ±1.0</td><td>7.2 ±0.8</td><td>16.0 ±4.5</td><td>28.8 ±6.8</td></tr></table>

We also used BLS measurements to study the magnetic damping in our samples. As mentioned above, the FMR measurement of the MLGr/MgO/CFB sample revealed two resonance peaks under an in-plane applied field. A detailed analysis requires careful deconvolution to extract the peak widths to investigate the magnetic damping and understand the different relaxation mechanisms in the studied stacks. To facilitate this, we used the full width half maximum (FWHM) of the BLS lines. We measured the FWHM of the Stokes and anti-Stokes lines and calculated their mean value $\delta F_{moy}$, which results from extrinsic and intrinsic contributions to damping. The experimental data were then fitted using equation (2) along the direction of the applied magnetic field, allowing extraction of the damping coefficient ($\alpha$) from the slope.

$$ \delta F_{moy} = 2\alpha \frac{\gamma}{2\pi} H + \delta F_{0} \quad (2) $$

$H$ is the in-plane applied field and $\delta F_0$ represents the linewidth that remains constant across magnetic fields, referred to as inhomogeneous broadening, arising from non-uniformities. Fig.3 exhibits the experimental data (symbols) and the fittings (continuous lines) obtained from Equation (2) for samples and controls. The deduced values of $\alpha$ are summarized in Table. 1.
<figure><img src="image_3.png" /><figcaption>Fig. 3. In-plane applied field-dependent BLS line-width of samples and controls.</figcaption></figure>

The linear fits applied to the line widths of the control samples (SiO₂/OxB/CFB) clearly show almost identical behaviour in the absence of MLGr. A significant increase in damping is observed in the presence of graphene. The enhancement is more marked for the MLGr/MgO/CFB sample.

Finally, we studied the SW non reciprocity to estimate the magnitude of iDMI [39,40] in our systems. As mentioned in the introduction section, iDMI is a necessary ingredient for stabilization of Néel-type domain walls (DWs) and for nucleation and propagation of topological textures such as skyrmions. Skyrmion-mediated switching of ferromagnets in presence of thermal perturbation has been demonstrated [26]. DMI has also shown to play a key role on STT switching in STT-MTJs [27–29]. Ab initio calculations [41] concluded, for a perfect stack, that DMI is an interface magnetic energy and is therefore very sensitive to the quality of the interfaces involving the ferromagnetic material. When it is present, iDMI leads, in BLS measurements, to a Stokes (Fₛ)/anti-Stokes (Fₛs) frequency difference (ΔF = Fₛ − Fₛs)
that gives direct information on the iDMI value using $\Delta F = \frac{2\gamma}{\pi M_s} D_{eff} k_{sw}$ where the effective iDMI constant $D_{eff}$ is related to the iDMI constant $D_s$ through $D_s = D_{eff} \times t_{FM}$.

Figure.4 exhibits the experimental spectra obtained, with an in-plane applied magnetic field of + 0.3 T and – 0.3 T, respectively, for the Gr based samples and controls, respectively. The spectra were recorded at a wave vector $k_{sw} = 20.45 \ \mu m^{-1}$, corresponding to an incidence angle of the laser beam of 60°; the maximum value available from our BLS set up, where the iDMI effect, if present, is the highest. Our results reveal that a vanishing iDMI is measured for all the samples.

<figure><img src="image_4.png" /><figcaption>**Fig. 4:** BLS spectra for structures with MLGr (left panel) and control ones, taken at $k = 20.45 \ \mu m^{-1}$, corresponding to an incidence angle of the laser beam of 60° ($k_{sw} = 20.45 \ \mu m^{-1}$) and an in-plane applied magnetic field of + 0.3 T and – 0.3 T. Symbols refer to experimental data and solid lines are Lorentzian fits for positive and negative field allowing comparison of Stokes and anti-Stokes frequencies.</figcaption></figure>
DISCUSSION

Raman results showed that, interestingly, the water-based ALD process (for Al₂O₃ deposition) used here seems to be less “destructive” to the MLGr, in contrast to the MBE process (for MgO deposition). Indeed, the MBE method led to an increase in graphene defects, a phenomenon supported by existing literature indicating potential non-uniformity and pinhole formation in films deposited through this method [17]. These defects could contribute to heterogeneity in the ferromagnetic layer, possibly arising from alterations in the CFB growth process. L. Zheng *et al.* reported pinhole-free and continuous coverage of graphene surface without any detected defects after Al₂O₃ deposition by water-based ALD [42]. F. Giannazzo *et al.* has also reported direct thermal water-based ALD of uniform Al₂O₃ thin films on MLGr preserving the graphene quality [43]. The static magnetic behaviour of our samples seems to indicate that the barriers attenuate the effects associated with the presence of graphene. Indeed, it should be noted that in the Si/SiO₂/Gr/FM systems, where the FM was directly deposited on Gr, a significant increase in *H* was reported compared to control samples without Gr [15,16,44]. This effect was for instance attributed to the presence of Gr ripples inducing additional heterogeneities in Gr/CoFeB heterostructures [16], or to increase magnetic disorder related to increase in defects in Gr/Co ones. In this latter case, the introduction of 1 nm MgO barrier leads to a strong attenuation of *H* enhancement, which is consistent with our observations. The increase of the in-plane saturation field indicates that the *presence of graphene* induces a more complex magnetization reversal mode (more domains), probably with greater magnetic heterogeneity, as observed for Gr/MgO/Co structures [15]. Our out-of-plane FMR measurements (supplementary material [36]) confirm such effect since the spectra show multiple modes concluding to the presence of areas with different well-defined properties compared to the control samples.
The BLS results indicate that the nature of the barrier do not affect significantly the effective magnetization $\mu_0 M_{eff}$. Indeed, in the absence of graphene, the values for MgO (0.97±0.01) T and Al₂O₃ (0.91±0.02) T are close and consequently also the effective perpendicular magnetic anisotropy constants $K_{eff}$, about (3.80±0.65)×10⁵ J/m³ (Al₂O₃) and (3.4±0.6)×10⁵ J/m³ (MgO), respectively. These results are consistent with those obtained from FMR measurements as shown in Table 1. In FM/oxide based-systems, an interfacial perpendicular anisotropy has been predicted by first-principles calculation and mainly attributed to hybridization of FM 3d and O 2p orbitals [45,46]. For instance, in MgO/CFB based-systems, it has been shown that the perpendicular anisotropy is entirely due to the CFB/MgO interfacial anisotropy [47]. Although it has been suggested that the PMA is mainly stabilized by the Fe-O hybridization at the CoFeB/MgO interface, the CoFeB/Ta interface has also been found to be crucial for the PMA in MgO/CoFeB/Ta structures [48,49]. More generally, the different crystalline states of the heavy metal (HM), diffusion of the boron atoms (in CoFeB) and the diffusion of HM at the CoFeB/MgO interface to affect the Fe 3d and O 2p orbital hybridization, were proposed as reasons for the HM's effect on the PMA properties [50–54]. First principles calculations of spin polarization and barrier-oxidation effects at the Co/Al₂O₃ interface in MTJs also indicate that the formation of Co–O interface bonds is crucial for mediating spin-polarized tunneling in Co/Al₂O₃ based junctions, via spin-polarized interface states [55].

By inserting a graphene monolayer (MLGr), through BLS results, we observe a decrease in $\mu_0 M_{eff}$, (0.67±0.01) T vs (0.91±0.02) T for Al₂O₃ barrier and (0.76±0.01) T vs 0.97±0.01) T for MgO barrier, and thus an increase in the effective perpendicular magnetic anisotropy constants $K_{eff}$, (5.2±0.6)×10⁵ J/m³ vs (3.80±0.65)×10⁵ J/m³ for Al₂O₃ barrier and (4.7±0.6)×10⁵ J/m³ vs (3.4±0.6)×10⁵ J/m³ for MgO barrier. This result highlights the ability of Gr to enhance PMA even when decoupled from the ferromagnetic layer (here CoFeB) by an oxide barrier. Similar behavior has been reported in the Gr/MgO/Co structures compared to
the MgO/Co ones [15], both capped with a 4 nm-thick MgO layer. It is worth noting that the decrease in µ₀Mₐff, and thus increase in PMA constant Kₐff, in the *presence of MLGr* is slightly more pronounced in the Al₂O₃-based system than in the MgO-based one. As discussed above, in such stacks, the oxidization rate, the oxygen diffusion at the oxide/FM interface, the diffusion of boron or Ta/FM intermixing can affect the effective magnetization and, consequently, the PMA. However, upon examination of the results reported in Fig.2(c), it becomes evident that *in the absence of graphene*, these effects, if present, seem to be similar for both systems. In fact, the control samples show virtually identical behavior, with their µ₀Mₐff and Kₐff almost indistinguishable, suggesting that the above mentioned effects cannot explain the observed magnetic properties. This difference in µ₀Mₐff values could be due to the specific deposition method employed. As outlined in the introduction, the MBE technique utilized for MgO deposition belongs to the PVD category, whereas the ALD technique employed for Al₂O₃ deposition is categorized under CVD method. It is possible that the growth mechanism of the two oxides on a substrate such as SiO₂ is different from the growth mechanism on a 2D material such as graphene here. It has been reported that Gr, as a substrate, potentially shapes the growth dynamics of the oxide layers in a heterostructure [56]. For instance, to obtain the formation of a smooth MgO overlayer, doping the graphene surface with Ti atoms, in the MBE deposition process of MgO, has been shown to mitigate surface atom mobility, enhance MgO adhesion, and reduce tunnel barrier roughness [21,57]. Compared to the FMR measurements, comparable result is found with the MLGr/Al₂O₃/CFB sample (µ₀Mₐff = 0.63±0.01) T, while the MLGr/MgO/CFB sample shows two resonance modes as evidenced in Fig.2(b). This results in two different values of µ₀Mₐff = (0.83±0.01) T and µ₀Mₐff = (0.68±0.02) T and extracted from the in-plane applied field dependent FMR frequencies shown in Fig.2(d), specific to the first and the second mode, respectively. The average µ₀Mₐff value is 0.75 T, which corresponds to the one obtained from BLS. This result
can be considered as consistent noting that that decoupled uniform modes are measured in
FMR and propagative modes in an average medium in BLS.

The results derived from the BLS line width analysis are presented in Table 1. The resulting damping coefficients exhibit remarkably close values *in the absence of the MLGr*, α = 7.6±1.0 ×10⁻³ for Al₂O₃ vs α = 7.2±0.8 ×10⁻³ for MgO, suggesting that both oxides have overall a comparable effect on magnetic damping. These values of α derived from our measurements are higher than the values of intrinsic damping coefficient documented in previous studies, such as αCFB = 4.2×10⁻³ for bulk Co₄₀Fe₄₀B₂₀ [58] and some CoFeB-based systems [59], αCFB = 4.9×10⁻³ for Co₄₀Fe₄₀B₂₀ (20 nm)/SiO₂ [60] but inferior for instance to αCo = 10.8×10⁻³ for Co (15 nm)/MgO [15]. In our analysis, the damping constant represents an effective value encompassing both intrinsic and extrinsic contributions (α = αCFB + αspin pumping + αTwo Magnon Scattering + αmosaicity). The additional damping probably likely originates from spin pumping contribution from the interface between the CoFeB and the Ta capping layers [59] and from additional heterogeneities due to Al₂O₃ and MgO sublayers leading to two-magnon and mosaicity contributions [57]. The *insertion of MLGr* was found to significantly increase damping constants α = 16.0±4.5 ×10⁻³ for Al₂O₃ *and* α = 28.8±6.8 ×10⁻³ for MgO barrier. δF₀ also increased for the two samples (MLGr/OxB/CFB) indicating the increase of heterogeneities in presence of MLGr. Indeed, defects and disorder in the FM layer give rise, as presented above, to two-magnon scattering and mosaic-like variation in the magnetocrystalline axes or other sample parameters [61]. In Gr/MgO/Co structures, compared to Co/MgO ones, all MgO capped, the observed enhancement of the FMR linewidth broadening in presence of graphene was attributed to effects such as two-magnon scattering or mosaicity, excluding spin pumping contribution [15]. A clear enhancement of the Gilbert damping was reported in permalloy (Py)/Cu/Gr stacks compared to that of Py or Py/Cu, even with 20 nm-thick Cu spacer. This resulted in a spin-mixing conductance surprisingly
comparable to the values obtained in Gr/Py samples [62]. However, the observed spin
pumping efficiency is difficult to reconcile with the expected backflow of spins according to
the standard spin pumping theory and the characteristics of graphene. Recently, it has been
reported that the Gr corrugations enhance damping of a Cu capped-CoFeB film through the
two-magnon contribution, concluding that MLGr layer is responsible for introducing
additional “roughness”, which is likely associated with the presence of graphene ripples. No
spin pumping was observed either without or with graphene [16]. Interestingly, the damping
coefficient of MLGr/Al₂O₃/CFB sample is much higher than that of Al₂O₃/CFB likely
associated to more inhomogeneities induced by the presence of graphene (ripples,
roughness,...). In the case of MgO spacer, the damping is also largely increased by the
graphene sublayer in line with the Raman measurements concluding to more defects. P. Sahu
et al. [63] observed in MgO/Gd-Alloyed BiₓSe₁-ₓ (BSG)/CoFeB/MgO/Ta heterostructures an
increase in interface roughness, regarding to controls without BSG, mainly attributed to the
BSG layer. High Resolution Transmission Electron Microscopy (HR TEM) analysis showed
that the CFB, MgO and Ta layers well follow the morphology of the BSG layer.
Quantitatively, it is worth noticing that our values obtained for the MLGr/OxB/CFB samples
are comparable to that (16.8×10⁻³) reported for the Co/MgO/Gr structures [15].

Finally, no significant frequency shift associated to SW non-reciprocity was measured
concluding that the studied structures (with and without graphene) present weak iDMI. The
contribution to the iDMI can originate from the MgO(Al₂O₃)/CoFeB and CoFeB/Ta
interfaces. Belabbes *et al.* by means of density functional theory (DFT) calculations showed
that the sign and magnitude of DMI can be entirely controlled by tuning the degree of
oxidation of the FM films and the associated 3d-2p hybridization. On another hand, DMI at
metallic 3d-5d interfaces is directly correlated to the 3d orbital occupation and the spin
flip/mixing processes with the spin-orbit active 5d states [64]. Due to its small spin orbit
coupling (SOC), Ta is expected to induce small iDMI at CoFeB/Ta interface [65]. For instance, it has been experimentally estimated by I. Mokhtari-Benguettat et al. [59] to be about 0.1 mJ/m², while R. Lo Conte et al. [66] reported a weaker value about 0.06 mJ/m². The MgO contribution to iDMI has been reported to be similar or slightly higher than that of Ta but of opposite sign: positive for Ta/CoFeB interface and negative for MgO/CoFeB one [59]. This could explain the weak (vanishing) iDMI measured in our samples.

The *presence of the MLGr* does not seem to change iDMI. This could mean that graphene has a predominantly geometric effect, but does not change the electronic configuration of the samples. It is to notice that the observed iDMI behavior can be desirable for STT-MTJs applications since it was reported that the DMI has a detrimental effect on STT switching in conventional heavy metal/ferromagnet single-interface STT [27,28].

## CONCLUSION

The aim of our study was to discriminate the role played by a monolayer graphene and the influence of oxide barriers, mainly on the magnetic properties of Gr/MgO (Al₂O₃) tunnel barriers/CoFeB heterostructures. For the control samples (*absence of Gr*), both static (VSM) and dynamic (MS-FMR and BLS) magnetic measurements showed that the type of oxide has no discernible effect on the coercivity, saturation field, effective magnetization, damping and interface Dzyloshinski-Moriya interaction (iDMI). Compared to the control samples, when *graphene is present*, while the coercivity appears to be little affected, the saturation occurs at higher fields compared to the controls indicating more complex magnetization reversal process. A decrease of effective magnetization is observed and thus an increase in the effective perpendicular magnetic anisotropy (PMA) constants, highlighting the ability of Gr to enhance PMA even when decoupled from the CoFeB layer by an oxide barrier. This effect is more pronounced in the Al₂O₃-based system than in the MgO-based one, probably a
consequence of the different deposition methods employed for the two barriers. A significant
increase of damping constants, more pronounced with MgO, is observed indicating the
increase of “heterogeneities” in presence of Gr giving rise likely to two-magnon and mosaic
contributions to the magnetic damping in our samples. Finally, the presence of the monolayer
graphene does not seem to change iDMI that was observed to be vanishing.

ACKNOWLEDGMENTS

ANR (Agence Nationale de la Recherche) and CGI (Commissariat à l'Investissement d'Avenir) are gratefully acknowledged for their partial financial support of this work through Labex SEAM (Science and Engineering for Advanced Materials and devices) ANR 11 LABX 086, ANR 11 IDEX 05 02 and ANR ACAF-20CE30-0027. This research was partially supported by structure fédérative de recherche NAP MOSAIC of the University Sorbonne Paris Nord through MONOGRAPH-IN and INTERFER-2D projects.

DATA AVAILABILITY

Data will be made available on reasonable request.

DECLARATION OF COMPETING INTEREST

The authors declare that they have no known competing financial interests or personal rela-
tionships that could have appeared to influence the work reported in this paper.
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