The magnetic Weyl semimetallic state in the chiral antiferromagnet Mn3Sn has attracted interest for its potential in memory technology. Despite vanishingly small magnetization, the material exhibits large transverse responses that can be electrically manipulated, similar to ferromagnets. Through deposition on heated Si/SiO2 substrates, we have fabricated polycrystalline Mn3Sn films that have coarse surfaces, the thinner of which have a discontinuous structure comprised of grains with diameters of the order of 100 nm. We confirm that these grains retain the anomalous Hall effect arising in the time reversal symmetry broken chiral antiferromagnetic phase of Mn3Sn at room temperature by serially connecting the grains with an additional conducting layer. These results pave the path for the potential applications of nanoscale Mn3Sn systems, which could be useful in the development of energy efficient memory devices.

Antiferromagnets (AFMs) have attracted significant attention as a promising material for next-generation high-integration-density and ultrafast information processing devices1–6 because of their magnetic field hardness,7 lack of perturbing stray field, absence of self-demagnetization,8 and terahertz-order spin dynamics.9,10 However, the vanishingly small net magnetization, which is the major origin for these benefits, also made it a challenge to control and detect the large electrical responses originating from their magnetic states. This issue often hinders the substitution of ferromagnets in spintronics devices with antiferromagnets.

On the other hand, the demonstration of electrical writing with (Néel) spin-orbit torque (SOT) and reading with relativistic responses such as anisotropic and spin Hall magnetoresistances in collinear antiferromagnets have broken through this stagnation.7,11–14 Almost simultaneously, another breakthrough occurred in non-collinear (chiral) antiferromagnets, where the time-reversal-symmetry (TRS) is macroscopically broken.15–23 As the first case in AFMs, the chiral antiferromagnet Mn3Sn has been experimentally found to exhibit magnetization-free macroscopic responses such as the anomalous Hall effect (AHE),15 the anomalous Nernst effect (ANE),24,25 and the magneto-optical Kerr effect (MOKE)26 due to its unique magnetic and electronic structures. In contrast to the relativistic responses, these take advantage of time-reversal operation, enabling to use the same writing/reading protocols and device structures27,28 as those developed for ferromagnets in conventional spintronics devices.29,30 Recently, magnetization-free responses originating from the macroscopic TRS have been also predicted in collinear antiferromagnets with unique crystal and magnetic symmetry.31 Such TRS-breaking antiferromagnets are being studied vigorously to develop functional antiferromagnets that can be used for practical devices.

Our target material, the noncollinear antiferromagnet Mn3Sn, has the hexagonal D019 structure with the ABAB stacking sequence of the Mn kagome planes along the (0001) direction [Fig. 1(a) and (b)]. This crystal geometry yields a three-sublattice anti-chiral antiferromagnetic spin structure below the Neel temperature TN420 K.32–34 This spin arrangement can be seen as a ferroic order of a cluster magnetic octupole. The octupole acts as a magnetic order parameter for the above-mentioned macroscopic responses requiring TRS breaking, similar to the magnetic dipole/magnetization in ferromagnets. In addition to this dominant order parameter, the spins can not slightly from the compensated arrangement to produce a small net ferromagnetic (FM) moment of ∼0.002 μB/Mn along the local easy axis (21¯1¯0⟩) in the (0001) plane.35 This canting moment corresponds to less than 0.1% of the local Mn moment. While this subdominant moment is essential for magnetic field control of the AF spin structure, it is not responsible for the AHE, ANE, and MOKE.15,24,35–37 Instead, it is the ferroic ordering of the cluster multipole that stabilizes the magnetic Weyl semimetallic states6,24,37,38 that give rise to a large Berry curvature in momentum space and induces AHE, ANE, and MOKE.15,24,35,36 Recent development in the techniques of the sample preparation and micro-fabrication has expanded the range of Mn3Sn's form from bulk single crystals to FIB-fabricated slabs and thin films, allowing us to probe responses to spin current39–41 and THz light42,43 such as the magnetic spin Hall effect and the THz Hall effect. Moreover, rapid progress is being made in pioneering device functions6 such as electrical control of the signal in a memory bit,27,28,44 its multi-level recording,8,27 and heat-flux sensing.8,45

The device miniaturization is essential for research and development of antiferromagnetic spintronics. Thus, it is very important to demonstrate that nano and micrometer scale clusters exhibit functional properties such as AHE without thermal fluctuations or device edge effects. Thus, the size and shape dependences of Mn3Sn are now being intensively studied.8,46,47 However, to date, nanometer to submicrometer sized samples have not been studied at room temperature.

In this work, we report the fabrication of isolated nano- and submicrometer-scale Mn3Sn grains through high-temperature deposition of Mn3Sn films immediately followed by cooling to room temperature. Although it is difficult to characterize the transport properties of the electrically as well as magnetically isolated grains, we have made it possible to measure the electrical properties by using a technique in which the grains are coated in series by a top electrode. As a result, we have confirmed the anomalous Hall effect, the characteristic function of the magnetic Weyl semimetallic state in Mn3Sn, in 100 nm-scale Mn3Sn grains at room temperature and at zero field. Our results pave the way for the application of an antiferromagnetic Weyl semimetallic state, i.e., a Weyl antiferromagnet in the spintronics devices (Fig. 1).

FIG. 1.

(a) Three-dimensional view and (b) two-dimensional view along the c-axis of the crystal and spin structures of Mn3Sn. The blue and orange spheres and arrows represent the Mn atoms and their spin, respectively, while the gray and black spheres represent the Sn atoms. Figures are generated with VESTA.62 (c) The XRD spectrum of a Mn3Sn film deposited on a Si/SiO2 substrate at 500 °C (t= 40 nm). The simulated diffraction pattern is shown on the bottom.

FIG. 1.

(a) Three-dimensional view and (b) two-dimensional view along the c-axis of the crystal and spin structures of Mn3Sn. The blue and orange spheres and arrows represent the Mn atoms and their spin, respectively, while the gray and black spheres represent the Sn atoms. Figures are generated with VESTA.62 (c) The XRD spectrum of a Mn3Sn film deposited on a Si/SiO2 substrate at 500 °C (t= 40 nm). The simulated diffraction pattern is shown on the bottom.

Close modal

For sample fabrication, we employ the DC magnetron sputtering method. Mn3Sn thin films are deposited at a high temperature of 500 °C on thermally oxidized Si substrates from a Mn2.7Sn target in a chamber with a base pressure of <5 × 10−7 Pa. The sputtering power and Ar pressure in the deposition process are 60 W and 0.7 Pa, respectively. The films are cooled to room temperature immediately after the deposition of the Mn3Sn film. This contrasts with previous works, where the Mn3Sn layers are annealed for some time after deposition.27,28,36,40,44,48–51 An Al capping layer (2 nm) is deposited on top of the Mn3Sn film in situ at room temperature to prevent oxidation. The composition of the Mn3Sn layers is determined to be Mn3.01(2)Sn0.99(2) by the scanning electron microscope energy-dispersive x-ray spectrometry (SEM-EDX). In the following, we refer to all our Mn-Sn films as Mn3Sn for clarity. The crystal structures of the films are investigated through x-ray diffraction (XRD) measurements. Standard four probe measurements are employed to measure both longitudinal and Hall resistivities of the films. These measurements are conducted using a variable temperature insert (VTI) equipped with a superconducting magnet (Teslatron PT, Oxford Instruments) applying a field in the out-of-plane direction swept at 0.13 T/min. The field dependence of the Hall resistivity is calculated by subtracting the longitudinal contribution, which is an even function with respect to the magnetic field and appears due to misalignment of electrical contacts for the measurements. The longitudinal resistivity ρxx and Hall resistivity ρyx are defined as VxIwtl and VyIt, respectively, where Vx and Vy are the voltages generated in the longitudinal and transverse directions under a current I in the longitudinal direction, w is the width of the sample, l is the length between the two electrodes that measure Vx, and t is the nominal thickness of the Mn3Sn layer of the film [Fig. 3(e)].

The structural properties are examined by XRD measurements at room temperature. We confirm that the films are the single phase of D019-Mn3Sn because the XRD spectrum of the Mn3Sn film (t = 40 nm) obtained by a 2θ/ω scan [Fig. 1(c)] shows the peaks expected only from D019-Mn3Sn and the Si/SiO2 substrate. Moreover, the ratio of the peak intensity is almost consistent with that of the simulation results for the randomly oriented D019-Mn3Sn. These results indicate that the film is polycrystalline with a mixture of crystallites with different orientations. The lattice constants are estimated to be a = 5.66 Å and c = 4.51 Å, which are almost the same as the previous report.36 

Before discussing the measurement results of the electrical transport, we first clarify the structural properties of the Mn3Sn films obtained by high-temperature deposition. Figures 2(a)–2(d) show the atomic force microscopy (AFM) images of the high-temperature-deposited films with varying t = 5–40 nm. As discussed below, two terminal sheet resistances R2w of the thicker films (t = 25–40 nm) are of several hundred Ω, characteristic to a metallic film. For the thinner films (t 20 nm), on the other hand, R2w is larger than several MΩ, i.e., the resistivity is too large for measuring the transport properties. In the following, we will describe such films (R2w > several MΩ) as “insulating” for simplicity. The AFM and cross-sectional transmission electron microscope (TEM) images confirm the isolated island structures of Mn3Sn in the t = 20 nm sample [Figs. 2(c) and 2(f)] in contrast to the continuous layer of Mn3Sn in the t = 25 nm sample [Figs. 2(b) and 2(e)]. The isolated grains have diameters as small as a hundred nanometers in the t = 20 nm sample. This length corresponds to several hundred times the lattice constants of Mn3Sn. The results obtained by these surface and cross-sectional images are consistent with the results of the above electrical measurements.

FIG. 2.

(a)–(d) AFM images of Mn3Sn films deposited on 500 °C Si/SiO2 substrates (t= 3–40 nm). Films with t25 nm are electrically conducting, while those with t20 nm are insulating. Between t=25 nm and t=20  nm, the gap between the submicrometer grains is formed, namely, islands of size ∼100 nm can be distinguished in the t=20 nm film, while no such isolated structures exist for the t=25 nm film; while troughs can be observed on the surface, the entire film seems to have a continuous structure. (e) and (f) Cross-sectional TEM images of films with t=25 nm (e) and t=20 nm (f). The structures of the films stipulated from the AFM images are confirmed: individual submicrometer structures are seen in the t=20 nm film, while no gaps can be observed in the t=25 nm film.

FIG. 2.

(a)–(d) AFM images of Mn3Sn films deposited on 500 °C Si/SiO2 substrates (t= 3–40 nm). Films with t25 nm are electrically conducting, while those with t20 nm are insulating. Between t=25 nm and t=20  nm, the gap between the submicrometer grains is formed, namely, islands of size ∼100 nm can be distinguished in the t=20 nm film, while no such isolated structures exist for the t=25 nm film; while troughs can be observed on the surface, the entire film seems to have a continuous structure. (e) and (f) Cross-sectional TEM images of films with t=25 nm (e) and t=20 nm (f). The structures of the films stipulated from the AFM images are confirmed: individual submicrometer structures are seen in the t=20 nm film, while no gaps can be observed in the t=25 nm film.

Close modal

As discussed earlier, the spontaneous anomalous Hall effect can be induced by the time reversal symmetry broken anti-chiral spin structure of Mn3Sn, which hosts the antiferromagnetic Weyl semimetallic state,38 and thus, it is a good indicator for determining the existence of the functional properties in the small grains. We carry out the magneto-transport measurements on a series of the Mn3Sn samples. Figure 3(a) shows the magnetic field dependence of the Hall resistivity ρyx(H) at 300 K for a film with t= 40 nm, where the field is applied perpendicular to the film. A clear hysteretic behavior of ρyx(H) with the zero-field component of ∼1.3 μΩcm and coercivity of ∼0.6 T is observed, similar to the values reported for post-annealed Mn3Sn thin films.8,36,48,52 On the other hand, the calculated Hall conductivity is ∼4 S/cm, several times smaller than the one for the post-annealed films due to the large longitudinal resistivity (ρxx = 559 μΩcm). This twice larger resistivity than that for the post-annealed sample36,48,52 should be attributed to the surface roughness and large number of grain boundary of the sample obtained in the high-temperature deposition process. As mentioned above, the formation of the isolated grains due to the high-temperature deposition is more pronounced for the thinner films. The t= 20 nm film shows insulating properties because of the gaps between the Mn3Sn grains. Such a grain structure is essential to confirm the functional properties including the anomalous Hall effect in the sub-100 nm grains. However, we cannot measure the magneto-transport properties in the electrically insulating sample.

FIG. 3.

(a) Magnetic field dependence of the Hall resistivity (AHE loop) of a film with t= 40 nm. The hysteresis loop is similar to those reported for post-annealed Mn3Sn thin films.8,36,48,52 (b) Magnetic field dependence of the Hall resistivity of a film with t = 20 nm. Note that to electrically connect the entire film, 2 nm of Al is deposited on the film in addition to the 2 nm Al capping layer used across all samples (4 nm total). The linear background of the Hall signal can be attributed to the ordinary Hall effect of the Al layer. (c) AHE loops of t= 40 nm and t= 20 nm samples normalized by their respective saturated values. The contribution of the ordinary Hall effect has been subtracted in both datasets. (d) Schematic of the serial connection achieved between the Mn3Sn grains. (e) Schematics of the Hall measurements.

FIG. 3.

(a) Magnetic field dependence of the Hall resistivity (AHE loop) of a film with t= 40 nm. The hysteresis loop is similar to those reported for post-annealed Mn3Sn thin films.8,36,48,52 (b) Magnetic field dependence of the Hall resistivity of a film with t = 20 nm. Note that to electrically connect the entire film, 2 nm of Al is deposited on the film in addition to the 2 nm Al capping layer used across all samples (4 nm total). The linear background of the Hall signal can be attributed to the ordinary Hall effect of the Al layer. (c) AHE loops of t= 40 nm and t= 20 nm samples normalized by their respective saturated values. The contribution of the ordinary Hall effect has been subtracted in both datasets. (d) Schematic of the serial connection achieved between the Mn3Sn grains. (e) Schematics of the Hall measurements.

Close modal

To probe the magneto-transport properties of the sub-100 nm Mn3Sn grains, we increase the thickness of the Al layer from 2 to 4 nm, thereby electrically bridging the gaps between the grains. Figure 3(b) shows ρyx(H) measured for such a film (t= 20 nm). In contrast to the film without the additional Al layer, the films exhibit the metallic transport behavior and a hysteretic behavior in ρyx(H). This Al conducting layer not only connects the isolated grains electrically in parallel but also causes a large shunting effect [Fig. 3(d)]. Therefore, the zero-field Hall resistivity is 0.06 μΩcm, which is two orders of magnitude smaller than that of the 40 nm film.

Assuming a 2 nm depth oxidation and bulk resistivity of 2.7 μΩ cm (Ref. 53) in the Al layer, we estimate that 95% of the current flows in the Al layer by using the two-resistor model described in Ref. 28. This corresponds to a Hall signal reduced by a factor of 20 compared to the case of a continuous t= 20 nm Mn3Sn film without any capping. As the measured Hall resistivity in the discontinuous film is smaller than the continuous 40 nm film by a factor of 21.7, it can be inferred that the intrinsic topological properties in the nanoscale grains are not significantly diminished. Furthermore, the field dependence of the normalized anomalous Hall response (calculated by dividing the Hall resistivities by their respective values at saturation), where the magneto-transport properties can be compared independently of the shunt effect, is identical to that of the t= 40 nm film [Fig. 3(c)]. Moreover, the temperature dependence of the remnant anomalous Hall effect is consistent with the one expected in the anti-chiral antiferromagnetic phase (inverse triangular spin structure phase) of Mn3Sn (Fig. S1).54,55 These results indicate that the 100 nm-scale Mn3Sn grains do, indeed, retain the spontaneous macroscopic responses originating from the ferroic order of cluster magnetic octupoles.

Finally, we estimate the magnetic anisotropy scale of our Mn3Sn particles based on their size and their measured magnetic stability. The energy barrier separating the stable magnetic states can be written as KV, where K is a constant and V is the volume of the particle, and its stability against thermal fluctuation can be described by the thermal stability factor ΔKVkBT (Ref. 56) (here, kB is the Boltzmann constant and T is the temperature). We approximate one of our Mn3Sn particles as a cylinder with the height of 20 nm and radius of 50 nm. One can estimate the lifetime of a stable magnetic configuration τ as τ=1f0eΔ, where f0 is an attempt frequency that models the number of times the magnetic orientation can thermally switch per unit time.57,58 In our case, let us adopt f0= 1 THz, as the attempt frequency should be on a similar order to the resonance frequency,59 and the latter of which has been measured to be ∼0.9 THz for the collective precession-like motion for Mn3Sn.60 When we take τ to be 1 min (corresponding to the time during the Hall measurement at ∼0 T), Δ32, which gives a lower bound of the relevant anisotropic energy coefficient as K800 J/m3 at 300 K. This figure is on the same order as the previously reported anisotropy energy coefficients for bulk Mn3Sn obtained by torque measurements under field.60,61

In conclusion, we have obtained isolated Mn3Sn grains with the size of ∼100 nm by using the high-temperature deposition on the thermally oxidized Si substrates. The anomalous Hall signal is clearly observed at room-temperature even in the isolated nanoscale Mn3Sn cells, potentially due to the interfacial anisotropy arising from the small system dimensions, and indicates possible functional responses even in smaller antiferromagnetic memory bits. Our findings give a first glance into the potential applications of nanoscale Mn3Sn systems, which could be useful in the development of energy efficient memory devices.

See the supplementary material for the temperature dependence of ρyx for the t= 20 nm film capped with 4 nm Al.

We thank S. Miwa for the support with the AFM measurements. This work was partially supported by JST-CREST (No. JPMJCR18T3), JST-MIRAI Program (No. JPMJMI20A1), and MEXT/JSPS-KAKENHI (Nos. 15H05882, 15H05883, 15K21732, and 19H00650). T.H. acknowledges support from the Hattori Hokokai Foundation. S.N. acknowledges support from the CIFAR as a Fellow of the CIFAR Quantum Materials Research Program. Institute for Quantum Matter, an Energy Frontier Research Center, was funded by DOE, Office of Science, Basic Energy Sciences under Award No. DE-SC0019331. The use of the facilities of the Materials Design and Characterization Laboratory at the Institute for Solid State Physics, the University of Tokyo, is gratefully acknowledged.

The authors have no conflicts to disclose.

Takumi Matsuo: Data curation (lead); Formal analysis (equal); Investigation (lead); Visualization (equal); Writing – original draft (lead); Writing – review and editing (equal). Tomoya Higo: Conceptualization (lead); Formal analysis (equal); Funding acquisition (supporting); Methodology (equal); Project administration (equal); Resources (equal); Supervision (equal); Visualization (equal); Writing – original draft (supporting); Writing – review and editing (equal). Daisuke Nishio-Hamane: Data curation (supporting); Writing – original draft (supporting); Writing – review and editing (supporting). Satoru Nakatsuji: Conceptualization (lead); Funding acquisition (lead); Project administration (lead); Resources (equal); Supervision (equal); Writing – original draft (supporting).

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

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