Compositional dependence of spintronic properties in Pt/GdCo films

Suzuki, Daniel H.; Lee, Byung Hun; Beach, Geoffrey S. D.

doi:10.1063/5.0165884

GdCo films have been widely used in spintronic applications, owing largely to their tunable degree of ferrimagnetic compensation. However, all key properties likewise depend on the alloy composition, and a systematic study of the interdependent spintronic properties with composition has not been reported. Here, we report the compositional dependence of key spintronic properties, including anisotropy, symmetric, and antisymmetric (Dzyaloshinskii–Moriya, DMI) exchange interactions, effective spin Hall angle, and domain wall mobility in a 3 nm Pt/GdCo composition series. We measure the magnetic anisotropy and determine an interfacial Pt/Co and bulk GdCo pair-ordering contribution to total anisotropy. Additionally, we estimate the exchange stiffness of all three interactions in GdCo as a function of composition. We conduct two types of domain wall motion experiments on patterned racetracks to determine the effective spin Hall angle and current-driven domain wall mobility. We find a 5× increase in effective spin Hall angle with increasing Gd concentration, suggesting an improvement in spin transfer efficiency in rare earth materials. Finally, we observe a monotonic decrease in the DMI strength with increasing Gd content, suggesting that DMI arises from the Pt/Co interfacial interaction.

Rare-earth (RE) transition-metal (TM) amorphous alloys are a promising class of ferrimagnetic materials for spintronics. RE-TM alloys consist of a RE sublattice antiferromagnetically coupled to the TM sublattice. Altering the proportion of RE and TM atoms allows one to control the magnetic properties of RE-TM films, including the saturation magnetization, $M_{s}$ ⁠, net spin density, coercivity, and magnetic anisotropy. This tunability has been used to engineer RE-TM films with low $M_{s}$ or low net angular momentum, in order to exhibit spintronic phenomena such ultra-small room temperature stable skyrmions,^1,2 high-speed current-driven domain wall motion,^2–5 and voltage-controlled skyrmion generation.⁶ However, despite the large volume of work performed on RE-TM films, a comprehensive study of the spintronic properties of RE-TM alloys, particularly in technologically relevant thin films, has yet to be conducted.

In this Letter, we analyze the magnetic properties of 3 nm GdCo across a wide range of compositions with the ultimate aim of quantifying the strength of the Dzyaloshinskii–Moriya interaction (DMI) as a function of composition using 1D domain wall (DW) motion experiments in patterned magnetic racetracks. DMI is an antisymmetric exchange interaction that stabilizes chiral spin textures, a necessity for skyrmion stability and current-driven DW motion. However, extracting this parameter from DW dynamics measurements requires knowledge of other parameters, such as DW width, which depend on material parameters that vary significantly with composition. Here, we undertake a systematic study to quantify all key parameters, providing a roadmap for compositional engineering of skyrmion host materials including magnetic anisotropy and exchange stiffness.

Films of Ta(4)/Pt(4)/Gd_xCo_1–x(3)/Ta(4)/Pt(2) were grown by D.C. magnetron sputter deposition (numbers in parenthesis indicate nominal thickness in nanometers) under a background pressure $< 2 \times 10^{- 7}$ Torr on thermally oxidized Si wafers with $0 \leq x \leq 0.7$ ⁠. The GdCo layer was grown by co-sputtering from Gd and Co targets with deposition rates of ∼3.5 and ∼1 nm/min, respectively, and with target power densities ranging from 0 to 8.8 and 4.9 W/cm², respectively. The lower Pt layer was selected to induce interfacial DMI and inject spin in the adjacent GdCo layer, enabling current-driven DW motion. A thin GdCo layer was chosen to maximize the interfacial DMI while minimizing potential bulk DMI contributions, as have been previously reported in thick RE-TM films.⁷ The Ta/Pt cap was used to prevent sample oxidation while preserving an asymmetric heterostructure for DMI maximization.

The 3 nm GdCo composition series saturation magnetizations, $M_{s}$ ⁠, are shown in Fig. 1(a). Magnetic compensation occurs at $x \approx 0.52$ ⁠, in contrast to the commonly reported bulk compensation composition of $x \approx 0.22$ ⁠.^8,9 We have previously explained this discrepancy in terms of a combined environment and dead layer model, in which a RE dead layer results in a reduced average magnetic moment per RE atom (RE atomic magnetic moment) in RE-TM films $< 10$ nm relative to thick films, and the average atomic magnetic moments of both the RE and TM decrease with increasing RE concentration at room temperature.¹⁰ In this model, the thickness of the RE dead layer is independent of RE-TM film thickness and increases with RE concentration, leading to expected Gd dead layer thicknesses of 0.5–2 nm for $x = 0.1 \to 0.7$ [see Fig. 4(d) of Suzuki et al.¹⁰]. Additionally, at room temperature, the two dominant exchange interactions in RE-TM films are the RE-TM and TM-TM interactions. This results in Gd and Co atomic magnetic moments that depend on the number of Co nearest neighbors, which decreases with increasing $x$ ⁠. Thus, at room temperature, the Gd and Co atomic magnetic moments decrease monotonically with increasing $x$ ⁠, approaching zero for pure Gd. The combined environment and dead layer model allows for the accurate estimation of individual sublattice contributions, $M_{G d}$ and $M_{C o}$ (where $M_{s} = |M_{G d} - M_{C o}|$ ⁠), in the grown series, shown as red and blue dashed lines in Fig. 1(a).

FIG. 1.

View large Download slide

Magnetic properties of 3 nm Gd_xCo_1−x films. (a) Saturation magnetization, $M_{s}$ ⁠, as a function of composition. The solid black line is the combined environment and dead layer model fit.¹⁰ The red and blue dashed lines are expected Gd and Co sublattice contributions predicted by the model (⁠ $M_{G d}$ and $M_{C o}$ ⁠), respectively. (b) Effective anisotropy energy density (black squares), $K_{u, eff}$ ⁠, and uniaxial energy density (red circles), $K_{u}$ ⁠, as a function of composition. The shaded region in (a) and (b) indicates perpendicular magnetic anisotropy.

Figure 1(b) shows the effective magnetic anisotropy, $K_{u, eff}$ ⁠, of the GdCo series, measured using hard-axis hysteresis loops via vibrating sample (VSM) and magneto-optical Kerr effect (MOKE) magnetometry. The films exhibit perpendicular magnetic anisotropy (PMA) for compositions $0.27 ≲ x ≲ 0.60$ ⁠. Subtracting the magnetostatic component gives the uniaxial anisotropy constant $K_{u} = K_{u, eff} + \frac{1}{2} μ_{0} M_{s}^{2}$ ⁠. $K_{u}$ decreases significantly with increasing Gd concentration at low $x$ and then increases and plateaus from $x = 0.25 \to 0.45$ before finally tending toward zero for $x > 0.5$ ⁠. The variation of $K_{u}$ with $x$ suggests two primary contributions. The most significant contribution is from the Pt/Co interfacial interaction, which is known to promote PMA in ultrathin Pt/Co films, overcoming the large magnetostatic penalty of the Co layer.^11–14 In GdCo, this interaction appears to drop rapidly with increasing Gd content. In addition, anisotropic pair–pair correlations have been found to introduce bulk PMA in RE-TM ferrimagnets.^15–19 This bulk PMA term increases with the number of RE-TM pairs, with pair models indicating a maximum in the pair-ordering induced anisotropy from $0.3 < x < 0.5$ ⁠, consistent with the observed $K_{u}$ trend in Fig. 1(b).

Before examining the DW dynamics, we consider the dependence of the exchange stiffness

A

with composition, which is necessary to predict the DW width

Δ = \sqrt{A / K_{u, eff}}

that plays a key role in the dynamical equations of motion. Symmetric exchange in GdCo arises from inter- and intra-sublattice contributions: two ferromagnetic interactions (Gd–Gd and Co–Co) and an antiferromagnetic Gd–Co interaction,

A_{tot} = A_{G d - G d} + A_{C o - C o} + A_{G d - C o}

⁠. The exchange stiffness of each interaction is given by^20,21

A_{i i} = \frac{1}{6} N_{a} S_{i}^{2} J_{i i} N_{c} r_{i i}^{2} x_{i}^{2}, A_{i j} = \frac{1}{3} N_{a} S_{i} S_{j} J_{i j} N_{c} r_{i j}^{2} x_{i} x_{j},

(1)

N_{a} = \frac{0.95}{x v_{G d} + (1 - x) v_{C o}},

(2)

with

S_{i}

the atomic spin,

J_{i j}

the exchange energy,

x_{i}

the atomic fraction,

N_{c} = 12

the coordination number,

r_{i j}

the interatomic spacing (⁠

r_{G d - G d} = 3.58 \times 10^{- 10} m, r_{G d - C o} = 3.02 \times 10^{- 10} m, r_{C o - C o} = 2.50 \times 10^{- 10} m

⁠),²⁰ and

N_{a}

the number of atoms per unit volume, assuming that the density of the amorphous film is 95% of its crystalline counterpart as has been reported previously.^21–23 The atomic volumes were taken to be

v_{G d} = 3.305 \times 10^{- 29} m^{3}

and

v_{C o} = 1.100 \times 10^{- 29} m^{3}

⁠, derived from their bulk densities. Thick film GdCo

J_{i j}

values from Hansen et al.⁹ and Gangulee and Kobliska²⁰ were used for calculations [Fig. 2(a)]. Both

J_{G d - G d}

and

J_{G d - C o}

are independent of composition; however,

J_{C o - C o}

decreases linearly with increasing Gd concentration. Additionally, in GdCo thin films, the Gd dead layer results in an alloyed composition that is Gd poor compared to its nominal composition. The true alloyed composition

\tilde{x}

can be estimated from the nominal composition

x

using the relations¹⁰^,

\tilde{x} = \frac{(t - t_{G d}) x v_{G d} - t_{G d} (1 - x) v_{C o}}{(t - t_{G d} x) v_{G d} - t_{G d} (1 - x) v_{C o}},

(3)

t_{G d} (x) = \frac{x v_{G d}}{x v_{G d} + (1 - x) v_{C o}} t_{G d, max},

(4)

with

t = 3

nm the nominal film thickness and

t_{G d, max} = 2.22

nm the maximum Gd dead layer thickness determined previously.¹⁰ Assuming

J_{C o - C o}

is dependent on

\tilde{x}

⁠, Eq. is used to determine

J_{C o - C o}

at nominal composition

x

⁠. To determine the

S_{i}

⁠, we utilize the combined environment and dead layer model referenced previously, which predicts the average atomic spin for each species in 3 nm GdCo [Fig. 2(b)]. The Gd spin value is reduced from its standard bulk value of

3.5 ℏ / a t

due to dilution with a non-magnetic Gd dead layer at room temperature and monotonically decreases with increasing

x

due to decreased Co nearest neighbors.

FIG. 2.

Symmetric exchange in GdCo films. (a) Previously reported exchange energy values in thick GdCo films as a function of alloyed Gd composition, x ̃ (see the text). Red and blue data points are from Hansen.9 The black dotted line is from Gangulee.20 (b) Expected average atomic spin of Gd and Co in 3 nm GdCo as a function of nominal composition, x. (c) Calculated 3 nm GdCo exchange stiffness, separated by each interaction. (d) Calculated domain wall width, Δ, of 3 nm GdCo samples exhibiting PMA.

View large Download slide

Symmetric exchange in GdCo films. (a) Previously reported exchange energy values in thick GdCo films as a function of alloyed Gd composition, $\tilde{x}$ (see the text). Red and blue data points are from Hansen.⁹ The black dotted line is from Gangulee.²⁰ (b) Expected average atomic spin of Gd and Co in 3 nm GdCo as a function of nominal composition, $x$ ⁠. (c) Calculated 3 nm GdCo exchange stiffness, separated by each interaction. (d) Calculated domain wall width, $Δ$ ⁠, of 3 nm GdCo samples exhibiting PMA.

Combining $J_{i j}$ and $S_{i}$ yields the exchange stiffness of each interaction in GdCo, shown in Fig. 2(c). As expected, due to the low Curie temperature of Gd (⁠ $T_{c, G d} \approx 290$ K), $A_{G d - G d}$ is insignificant compared to the other two interactions at room temperature. Additionally, the Co–Co interaction dominates except at relatively high $x > 0.4$ ⁠. The value of $A_{tot}$ is reduced by an order of magnitude from $x = 0 \to 0.5$ as the Co–Co exchange is reduced due to fewer Co–Co pairs. We have estimated $Δ = \sqrt{A_{tot} / K_{u, eff}}$ using the sum of the inter- and intra-sublattice exchange stiffness, as shown in Fig. 2(d). Despite the large decrease in exchange stiffness over this range, the corresponding reduction in anisotropy results in a relatively weak variation of $Δ$ ⁠, ranging from ∼3 to 9 nm in the studied range.

GdCo PMA samples with

x = 0.27 \to 0.49

were patterned into

40 \times 100

μm² magnetic racetracks for DW motion experiments [Fig. 3(a)]. Kerr microscopy was used to observe DW displacement. Two types of experiments were conducted: depinning and velocity measurements. In the depinning experiment, a low DC density,

j (\sim 10^{9} A / m^{2})

⁠, was applied along the length of the track and an assisting out-of-plane (OOP) magnetic field was slowly ramped from zero until a critical field,

H_{d p}

⁠, was reached and the propagated freely along a length of the track of at least

\sim 30 μ m

⁠. With no applied longitudinal field,

H_{x}

⁠, both up–down and down–up DWs exhibit the same linear relationship between

j

and

H_{d p}

⁠, consistent with previous observations [Fig. 3(b)].^2,24 The slope of Fig. 3(b) gives the spin–orbit torque (SOT) effective field per unit current density,

χ

⁠, the proportionality constant between current density

j

and spin Hall effective field,

H_{S H} = χ j

⁠. Figure 3(c) shows

χ

as a function of composition.

χ

is found to exponentially increase with decreasing

M_{s}

[see Fig. 1(a)], consistent with previous reports near magnetic compensation.^25–27 To separate the compositional dependence of

χ

from its

M_{s}

dependence, the effective spin-Hall angle was calculated using²⁸^,

θ_{S H}^{eff} = \frac{4 e μ_{0} M_{s} t}{π ℏ} χ,

(5)

where

e

is elementary charge.

θ_{S H}^{eff}

is found to increase with increasing

x

[Fig. 3(d)]. A similar increase in

θ_{S H}^{eff}

in TbCo with increasing Tb content was reported in Ref. 29, and theoretical work has shown significant SOT generated by RE atoms, particularly when alloyed.³⁰ The determination of

θ_{S H}^{eff}

here relies on DW depinning measurements that observe net SOT phenomena. As a result, the extracted

θ_{S H}^{eff}

does not distinguish between individual sublattices (that is, Pt/Co vs Pt/Gd interaction); only the combined effect is measured. The increase in

θ_{S H}^{eff}

with RE content suggests that the RE sublattice enhances spin transmission in RE-TM alloys.

FIG. 3.

View large Download slide

Domain wall depinning experiments. (a) Image of magnetic racetrack device used for de-pinning and DW velocity experiments showing a down-up DW. Inset, series of images showing DW displacement in Gd_0.47Co_0.53 with $j_{drive} = + 11.5 \times 10^{11}$ A/m² and a train of 5 ns current pulses totaling 100 ns between images. (b) Exemplary plot of depinning field, $H_{d p}$ ⁠, vs applied DC drive current, $j$ ⁠, in Gd_0.49Co_0.51. (c) SOT effective field per unit current density of 3 nm GdCo films as a function of composition. (d) Effective spin-Hall angle as a function of composition.

We next examine DW velocity

v

as a function of

j

and

H_{x}

⁠. DW displacement was measured by Kerr microscopy. Current pulses of 5 ns were used to drive DW motion. In magnetic racetracks, the DMI stabilizes homochiral Néel DWs via an effective field,

H_{D}

⁠, perpendicular to the wall. The homochiral Néel walls enable translational current-induced DW motion. The DW velocity,

v

⁠, in ferrimagnetic materials is well-described by the 1D model:³¹^,

v = \frac{\frac{π}{2} γ_{eff} Δ μ_{0} H_{S H}}{\sqrt{α_{eff}^{2} + {(\frac{H_{S H}}{H_{D} + H_{x}})}^{2}}},

(6)

where

γ_{eff}

is the effective gyromagnetic ratio and

α_{eff}

is the effective Gilbert damping, ferrimagnetic analogues of the ferromagnetic magnetic properties.^2,32

v

was first measured under

H_{x} = 0

and varying

j

[Figs. 4(a)–4(c)].

v

initially increases linearly with

j

before saturating at an asymptotic velocity,

v_{max} \approx \frac{π}{2} γ_{eff} Δ μ_{0} (H_{D} + H_{x})

⁠. In the saturation regime,

v

varies linearly with

H_{x}

⁠, and extrapolation to zero yields

H_{D}

⁠. However, the saturation current density,

j_{sat}

⁠, which we define here as the current density required to reach 90% of

v_{max}

with

H_{x} = 0

⁠, is composition dependent. This is evidenced by comparison of Figs. 4(a) and 4(b), in which

v

approaches saturation at

j_{sat} \approx 16 \times 10^{11}

A/m² and

j_{sat} > 20 \times 10^{11}

A/m², respectively. Figures 4(d)–4(f) show measurements of

v

vs

H_{x}

at three representative compositions. The DMI-stabilized homochiral DWs result in

H_{D}

with opposite signs for up–down and down–up DWs. As a consequence, up-down and down-up DWs have inverse responses to

H_{x}

⁠, i.e., larger

H_{x}

decreases

v_{u p - down}

but increases

v_{down - u p}

[Figs. 4(d)–4(f)]. The consistent sign of

H_{D}

for each type of DW across

x = 0.27 \to 0.49

indicates no chirality change (or equivalently sign flip in

D

⁠) across this composition range. As

j

decreases below

j_{sat}

⁠, the

v

vs

H_{x}

data become increasingly nonlinear [Fig. 4(e)]. A linear fit assuming saturation behavior would result in a large overestimation of

H_{D}

⁠. To properly account for the curvature of

v

vs

H_{x}

at

v < v_{max}

⁠, the data in Figs. 4(a)–4(c) and Figs. 4(d)–4(f) were fitted to the full 1D model [Eq. ]. The

χ

of each sample determined from the previous depinning experiments [Fig. 3(c)] was used to convert the applied current density to a spin-Hall effective field. The fits of the 1D model are shown as solid lines in Fig. 4.

FIG. 4.

View large Download slide

Current-induced domain wall velocity measurements in Pt/GdCo heterostructures at three representative compositions. (a)–(c) Domain wall velocity, $v$ ⁠, vs applied current density, $j$ ⁠, with no applied longitudinal field. Dotted line indicates $j$ used for corresponding $v$ vs $H_{x}$ measurement below. (d)–(f) Domain wall velocity vs applied longitudinal magnetic field, $H_{x}$ ⁠. Solid lines are fits to the 1D model (see the text). The error bars are the standard error of 5–20 measurements.

The combined results of the DW motion experiments are shown in Fig. 5. Both

α_{eff}

and

g_{eff}

[Figs. 5(a) and 5(d)] show divergent behavior near

x \approx 0.42

⁠, suggesting that the angular momentum compensation composition of 3 nm GdCo is near that point. We observe

> 10 \times

higher

α_{eff}

on the Gd-dominated side of angular momentum compensation compared the Co-dominated side. We attribute this significant increase in

α_{eff}

to the presence of the Gd dead layer potentially increasing surface roughness, which has been shown to enhance damping.³³

α_{eff}

and

γ_{eff}

are combined to determine the DW mobility

μ

⁠,

μ = \frac{π γ_{eff} Δ}{2 α_{eff}}

(7)

shown in Fig. 5(e).

μ

decreases exponentially with increasing Gd content, primarily due to the reduction in the DW width,

Δ

⁠, over this range. The ∼30x reduction in

μ

from

x = 0.27 \to 0.49

is almost equivalent to the increase in

χ

[Fig. 3(c)] observed over the same range, resulting in near-constant slopes in the

v

vs

j

data at all

x

at small

j

[Figs. 4(a)–4(c)].

FIG. 5.

View large Download slide

Pt/Gd_xCo_1−x dynamic properties as a function of composition. (a) Effective damping. (b) Maximum domain wall velocity. (c) DMI effective field. (d) Effective g-factor. (e) Domain wall mobility. (f) DMI magnitude.

The saturation velocity, $v_{max}$ ⁠, at $H_{x} = 0$ is shown in Fig. 5(b). With no applied longitudinal field, $v_{max} = \frac{π D}{2 S}$ ⁠, with $S = \frac{ℏ M_{s}}{μ_{B} g_{eff}}$ the net spin density. The $v_{max}$ data exhibit constant $v_{max} \approx 300$ m/s except near the angular momentum compensation point (⁠ $S = 0$ ⁠), where $v_{max} \approx 400$ m/s. This is consistent with previous investigations of ferrimagnetic materials near angular momentum compensation where the DW velocity has similarly increased with reduced spin density.^2–4 The $v_{max}$ data suggest that except at $S \approx 0$ ⁠, the ratio $D / S$ remains approximately constant.

Finally, the effective DMI field,

H_{D}

⁠, was extracted from the

v

vs

H_{x}

data and is presented in Fig. 5(c). The measured

H_{D}

ranges from 250 to ∼1700 Oe with a minimum at

x = 0.4

⁠. There is no clear

H_{D}

trend evident as a function of composition. However,

H_{D}

depends on

M_{s}

and

Δ

⁠, which both also depend on composition. To isolate the strength of the DMI as a function of Gd composition,

H_{D}

was converted to

D

using^31,34

D = μ_{0} M_{s} H_{D} Δ .

(8)

The results are plotted in Fig. 5(f), where a clear trend emerges. We observe a monotonic decrease with increasing Gd content across the entire PMA composition range, suggesting that the Pt/Co interaction is primarily responsible for the occurrence of DMI in GdCo films and that the addition of Gd only dilutes the strength of the DMI. The value of

D

at

x = 0.44

is smaller than the reported DMI in a 6 nm film at the same composition.² At

x = 0.27

⁠, we find a similar value of

D

compared to a previous report on a 5 nm film.³⁵ In heavy metal/ferromagnetic systems,

D \propto \frac{1}{t}

when induced at an interface,³⁶ making the reduced

D

values observed in 3 nm GdCo surprising. Composition gradients in RE-TM alloys have been shown to induce a bulk DMI with

D \propto t

⁠.⁷ This could add a second contribution to the DMI in films

> 3

nm that results in larger

D

values that is insignificant in our thinner films. We note that the measured

χ

values [Fig. 3(c)] are also lower than the reported SOT in 6 nm GdCo,² which, combined with a reduced

D

⁠, results in lower DW velocities than previously observed. Because DMI is sensitive to interfacial quality,^37–40 we ascribe both of these effects to increased surface roughness in thin GdCo films, potentially due to the significant dead fraction of Gd present in 3 nm films, particularly at high Gd concentrations.

In summary, we have analyzed the spintronic properties of 3 nm GdCo. A number of GdCo magnetic properties critical to spintronic devices and stabilization of chiral spin textures are found to strongly depend on RE concentration. The total effective anisotropy is found to strongly decrease with increased Gd concentration, consistent with a Pt/Co interfacial origin. We also observe a sub-dominant increase in uniaxial anisotropy that is maximized at $x \approx 0.4$ ⁠, consistent with a pair-ordering bulk anisotropy origin that is often used to explain bulk PMA in RE-TM alloys. The combination of DW velocity and depinning measurements under longitudinal and polar applied fields allows for determination of a variety of dynamic magnetic properties in Pt/GdCo heterostructures, including DMI effective field, damping, spin density, spin transfer efficiency, and effective spin-Hall angle. These fundamental measurements are combined with knowledge of the ferromagnetic and antiferromagnetic exchange interactions in GdCo to determine the magnitude of antisymmetric exchange constant $D$ ⁠. $D$ is found to monotonically decrease with increasing Gd content, suggesting that the Pt/Co interface is the primary source of DMI in Pt/GdCo films. This thorough characterization of GdCo's spintronic properties should allow for improved design of materials for next-generation DW and skyrmionics devices.

This work was supported in part by the DARPA TEE Program and the Samsung Global Research Outreach (GRO) Program.

AUTHOR DECLARATIONS

Conflict of Interest

The authors have no conflicts to disclose.

Author Contributions

Daniel Hiroshi Suzuki: Conceptualization (equal); Data curation (equal); Formal analysis (equal); Investigation (equal); Methodology (equal); Software (equal); Visualization (equal); Writing – original draft (equal); Writing – review & editing (equal). Byung Hun Lee: Data curation (supporting); Formal analysis (supporting); Methodology (supporting); Writing – review & editing (supporting). Geoffrey S. D. Beach: Conceptualization (equal); Formal analysis (equal); Funding acquisition (equal); Methodology (equal); Project administration (equal); Resources (equal); Supervision (equal); Writing – review & editing (equal).

DATA AVAILABILITY

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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Compositional dependence of spintronic properties in Pt/GdCo films

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Compositional dependence of spintronic properties in Pt/GdCo films

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