Superlattices of antiferromagnetic μ-Fe2O3 and diamagnetic β-Ga2O3 are grown by plasma-assisted molecular beam epitaxy on (010) oriented β-Ga2O3 substrates in which ferromagnetism emerges above room temperature. To investigate the suspected interface origin of the ferromagnetic phase, identical superlattice structures are grown at various substrate temperatures and beam fluxes. Atomic-resolution scanning transmission electron microscopy images confirm the registry of μ-Fe2O3 to the β-Ga2O3 layers in these superlattices. Atomic force microscopy and high-resolution x-ray diffraction are used to examine the growth morphology and characterize the superlattice interface roughness. The saturation magnetization of the ferromagnetic phase is observed to increase strongly with the interface roughness. Conversely, smoother superlattices exhibit a weaker ferromagnetic response and a higher density of paramagnetic moments along with evidence of superparamagnetic clusters. These findings are consistent with the interface origin for the ferromagnetic response in these superlattices. The demonstration of an interface magnetic phase in nearly lattice-matched monoclinic Fe2O3/Ga2O3 opens the door to ultrawide bandgap heterostructure-engineered magnetoelectronic devices, where ferromagnetic switching of the interface phase can be incorporated into high-field devices.

A new monoclinic μ-phase of Fe2O3, isomorphic to β-Ga2O3, was recently stabilized using plasma-assisted molecular beam epitaxy (PAMBE).1 Scanning transmission electron microscopy (STEM) structural and chemical imaging showed the registry of the μ-Fe2O3 to the β-Ga2O3 lattice, indicating that high-quality epitaxial growth of these heterostructures was feasible. While bulk μ-Fe2O3 was predicted to be antiferromagnetic, the μ-Fe2O3/β-Ga2O3 superlattices show ferromagnetic hysteresis at room temperature.1 

β-Ga2O3 is a wide bang gap (∼4.6 eV) semiconductor with a very high breakdown voltage (Eb = 8 MV/cm) being developed for high-power devices.2,3 The monoclinic structure of β-Ga2O3 results in anisotropic properties with the growth plane (010) having provided the highest growth rate compared to any other orientation, such as (2¯01) and (100).4 The electron spin lifetimes (on the order of 100 ns)5 and the high breakdown voltage2 make β-Ga2O3 a good candidate for spin-based field effect transistors. These devices typically require a magnetic semiconductor component to provide spin-split valence and/or conduction band edges, which could be achieved by alloying or heterostructuring β-Ga2O3 with Fe3+-rich (spin-5/2) layers. β-Ga2O3 also has a low lattice thermal conductivity,6 which is actually beneficial for thermoelectric power factors. While wide bandgap semiconductors are generally bad candidates for thermoelectrics, magnon-electron drag in combination with low lattice thermal conductivity might enable large thermoelectric power factors.7 Previously, ferromagnetism at room temperature was reported in Mn2+ and Fe3+ doped β-Ga2O3.8,9 Density functional theory (DFT) calculations have suggested that Ga cation vacancies at the tetrahedral and octahedral sites result in spin polarization due to the local moments of the O2− anions.10 Gallium vacancy defects can have charge states from 0 to 3−.11 They are paramagnetic in the 2−, 1−, and 0 charge states, and are diamagnetic in the 3− charge state.11 However, the β-Ga2O3 substrates used in both this study and the study from Jamison et al.1 were found to be purely diamagnetic.

Previous DFT calculations predicted that μ-Fe2O3 should exhibit antiferromagnetic order with Fe3+ spins on equal numbers of octahedral and tetrahedral sites coupled antiferromagnetically, presumably through the superexchange mechanism.11 The theoretical lattice constants for μ-Fe2O3 from the DFT calculations were within 1.5% of those of β-Ga2O3, resulting in a strain energy of 0.03 eV per formula unit.1 According to the DFT calculations, ferromagnetism emerges at the interface region between β-Ga2O3 and μ-Fe2O3 layers where Ga3+ and Fe3+ occupy alternating tetrahedral and octahedral sites.1 

Here, we test the interfacial magnetism hypothesis from Jamison et al. by investigating how variations in the roughness of the interface layer between β-Ga2O3 and μ-Fe2O3 layers affect the observed magnetic response. Rougher epilayers exhibit higher saturation magnetizations consistent with higher fraction of Fe3+ in the ferromagnetic state.

μ-Fe2O3/β-Ga2O3 superlattices are grown via PAMBE in a Riber/MBE Control M7 system equipped with Ga and Fe effusion cells and a Veeco oxygen plasma source with a forward power of 300 W (with ≤1 W reflected). The flow of O2 is continuously monitored and maintained at a growth chamber pressure of 1.5 × 10−5 Torr. The flux of Fe is held constant between samples with a beam equivalent pressure of 3.4 × 10−8 Torr. Each μ-Fe2O3 layer is deposited for 30 s by cycling the Fe-shutter. The substrate temperature Tsub (as measured through the thermocouple) is varied between 500 and 700 °C. β-Ga2O3 spacers are grown by varying the Ga flux (ΦGa) between 2.5 and 6 × 10−8 Torr with a constant deposition time at 400 s.

The substrates used for these growths are 5 × 5 × 0.5 mm3 unintentionally doped (010) β-Ga2O3 single crystals (Tamura Corporation). The substrates are cleaned by sonication for 5 min in acetone followed by the same sonication procedure in methanol and isopropanol. Cleaned substrates are indium-bonded to the unpolished side of a 3 in. silicon wafer. A final cleaning is performed in the growth chamber by heating the substrates to 800 °C for 10 min while being exposed to the O2 plasma. Following the pregrowth cleaning procedure, the substrates are cooled to their final growth temperatures.

The layer structure is designed to investigate how variations in the interface quality between μ-Fe2O3 and β-Ga2O3 affect the ferromagnetic response. First, a 100-nm-thick buffer layer of β-Ga2O3 is grown followed by a superlattice with a constant thickness of μ-Fe2O3 (∼3 ml) with ∼8–16 nm-thick β-Ga2O3 spacers, depending on the Ga flux. A total of 10 periods are grown with the final layer being a capping layer of β-Ga2O3.

Figure 1(a) shows a cross-sectional STEM high-angle annular dark field (HAADF) image of a superlattice grown at Tsub = 700 °C and ΦGa = 6 × 10−8 Torr. STEM data are recorded using a probe corrected Themis-Z S/TEM instrument operated at 200 kV for the HAADF image and an FEI Titan 60/300 STEM instrument operated at 300 kV with a beam current of 60–90 pA for the energy dispersive x-ray (EDX) elemental map. Contrast in the HAADF image is related to the atomic number; thus, the Fe oxide layer exhibits dark contrast relative to the β-Ga2O3 spacers. The EDX elemental map in Fig. 1(b) confirms the identification of the layers in the superlattice and Au on the top protective layer. The HAADF-STEM image at higher magnification [Fig. 1(c)] shows the μ-Fe2O3/β-Ga2O3 interfaces along the [001] direction. While the epitaxial relationship is well defined and there are no detectable defects at the interfaces, the interfaces between μ-Fe2O3 and β-Ga2O3 are undulating [Fig. 1(c)]. This may be associated with roughness due to the island-layer-by-layer growth mode or elemental diffusion between Fe and Ga around their interfaces.

The entire growth (i.e., during both Ga and Fe deposition) is observed by in situ reflection high energy electron diffraction (RHEED), representative patterns shown in Figs. 1(d) and 1(e) indicative of 2D growth without an identifiable change in lattice parameters or surface structure.

To investigate the epilayer morphology of the superlattices, atomic force microscopy (AFM) is performed using the Bruker AXS Dimension Icon. Figure 2 shows 1 × 1 μm2 scans of the starting β-Ga2O3 substrate and different superlattices grown at various substrate temperatures. The bare (010) β-Ga2O3 substrate has a surface roughness of 0.205 nm [∼1 monolayer (ML) roughness].

The AFM surface morphologies of the superlattices appear similar to previously reported β-Ga2O3 (010) PAMBE growths at these temperatures.12 Islandlike plateaus are clearly seen in the AFM data for all the samples, regardless of the growth conditions. This is indicative of island-layer-by-layer growth or the Frank–van der Merwe (FM) growth mode and is consistent with the observed 2D RHEED patterns. The FM growth mode is only stabilized for the so-called “strong” substrates whose lattice constant is imposed on to the film atoms,13 which implies near-perfect lattice matching between the substrate and the growth layers. Because the theoretical lattice constants of μ-Fe2O3 are within ∼1.5% β-Ga2O3,1 the observations of the FM growth mode seen in the AFM data confirm that the μ-Fe2O3 layers are lattice-matched to β-Ga2O3 across the 10 period superlattice structure.

The surface roughness of the superlattices is strongly correlated with the substrate temperature Tsub [Fig. 2(e)]. Samples grown at Tsub = 700 °C have surface roughness measurements close to ∼0.5 nm (2 ML), whereas those grown at Tsub = 500 °C have large surface roughness close to ∼1.1 nm (4 ML). This suggests that the SL interface quality deteriorates at low Tsub perhaps due to statistical roughening caused by a reduced adatom diffusion length. The surface roughness measurements are also correlated with ΦGa. Superlattices grown at ΦGa = 6 × 10−8 Torr have smoother surfaces (RMS ∼ 0.54 nm), and thus have better interface quality than those grown at ΦGa = 2.5 × 10−8 Torr (RMS ∼0.70 nm). The smoothest surfaces are observed at high ΦGa and high Tsub.

X-ray diffraction (XRD) is performed using a Bede diffractometer in the triple axis mode and a 4-bounce (220) Ge monochromator. High-resolution XRD in Fig. 3 shows the presence of superlattice peaks (SL0, SLn+1) alongside the substrate peak for all the samples in this study. Jamison et al. determined the b lattice parameter for μ-Fe2O3 through CADEM (Ref. 14) XRD fitting of the superlattice XRD data as 3.12 ± 0.4 Å, consistent with DFT calculations.1 The average b lattice parameter for the superlattices is calculated from the interplanar spacing (d020) of the (020) 0th order SL0 diffraction peak assuming a monoclinic structure for the superlattices and ranges from 3.046 to 3.055 Å. The individual Fe2O3 layer thicknesses range from 2.86 to 4.11 ml. The average Ga2O3 spacer thickness is ∼29 ml at ΦGa = 2.5 × 10−8 Torr and ∼53 ml at ΦGa = 6.0 × 10−8 Torr. Laue oscillations are noted in the XRD data but disappear at low Tsub and low ΦGa, indicative of higher interface roughness. The SL−1 peaks in the XRD data are narrower at high Tsub and high ΦGa (large spacer thickness) (Fig. 3). The full width at half maximum (FWHM) of the SL−1 peak is strongly anti-correlated with ΦGa [Fig. 3(b)]. The FWHM at ΦGa = 6 × 10−8 Torr is ∼ 2 × smaller than at ΦGa = 2.5 × 10−8 Torr. The FWHM of the XRD satellite peaks are indicative of the coherency of the SL structure and are sensitive not just to interface roughness but also to strain inhomogeneity, which can arise due to composition fluctuations or partial strain relaxation.

Superconducting quantum interference device (SQUID) magnetometry is performed on the SL structures to investigate their magnetic properties. The applied field (H) is scanned in a hysteretic fashion between ±80 kOe along the [102] direction (in plane) and the induced moment is measured. The diamagnetic background is removed by measuring magnetization versus H of a bare β-Ga2O3 substrate, which exhibits a purely diamagnetic response. The magnetization (M) of the μ-Fe2O3 layers is obtained by normalizing the magnetic moment of each sample by the volume of μ-Fe2O3 within each sample as determined by the XRD measurements. Ferromagnetic hysteresis (non-zero coercivity and remanence) is observed across the measured temperature range (5–300 K) as shown in Fig. 4(a). A paramagnetic component emerges at low temperature.

Isolated (dilute) Fe3+ ions should exhibit a paramagnetic response; thus, the diffusion of Fe3+ into the β-Ga2O3 spacers is likely the reason for the paramagnetic component. Assuming that at high-field the ferromagnetic component is saturated, and the antiferromagnetic component has negligible magnetization, we fit the high-field (|H| > 10 kOe) magnetization data to the Brillouin function [Eqs. (1) and (2)],

x=gμBJHk(Tθ),
(1)
M=NpgμBJ[((2J+1)2J)coth((2J+1)x2J)12Jcoth(x2J)]+Msat,
(2)

where θ is the paramagnetic Curie temperature, Np is the density (cm−3) of paramagnetic moments, g = 2.02 is the g-factor, J = 5/2 for Fe3+, T is the measurement temperature, and Msat is the saturation magnetization of the ferromagnetic component. Examples of M versus T data at T<40K are shown in Fig. 5(a), which are fit to the temperature dependent Brillouin function using just two free parameters, θ and Np. Next, the M versus H data for this sample, taken at T = 300 K [Fig. 5(b)], are fit to the field-dependent Brillouin function using a single parameter (Msat) assuming that the previously acquired values of θ and Np do not vary with the temperature. Following this procedure, the paramagnetic component at all temperatures and fields is described by Eq. (2) and is subtracted from the data to isolate the purely ferromagnetic response [Fig. 4(c)]. The median paramagnetic Curie temperature for these superlattices is θ=3.55±1.03K, which indicates an antiferromagnetic coupling between dilute Fe3+ ions. Similar behavior was previously studied in II–VI dilute magnetic semiconductors in the low alloying regime and arises from small clusters of antiferromagnetically coupled spins via short-range superexchange.15 The magnetic properties of the superlattices, including the Brillouin parameters, are summarized in Table I.

The M versus T behavior is obtained in the standard zero-field-cooled (ZFC) and field-cooled (FC) conditions using a SQUID magnetometer with an applied field of 500 Oe. Figure 6(a) shows ZFC and FC curves for a superlattice grown at Tsub = 700 °C and ΦGa = 6 × 10−8 Torr. The discrepancy between the ZFC and FC data is indicative of superparamagnetic clusters (SCs), i.e., a distribution of single-domain ferromagnetic clusters, whose spontaneous magnetization averages to zero over time due to thermal fluctuations.16 As described by Bruvera et al., the peak at T ∼ 20.5 K in the ZFC curve does not correspond with a representative blocking temperature (TB) for the SCs, but rather is governed by the details of the cluster size distribution.17 A more reliable method involves taking the derivative with respect to the temperature of the difference between the ZFC-FC data, d(ZFC-FC)/dT [Fig. 6(b)], which represents the distribution of TB18 associated with the size distribution of SCs. The average TB is estimated by fitting a Gaussian to the d(ZFC-FC)/dT data (TB = 10.1 K for the sample in Fig. 6). Superparamagnetic response is observed in samples grown at Tsub = 700 °C with TB ranging from 9.8 to 10.7 K. At Tsub = 600 °C, this behavior is less evident with TB = 13.1 K. For superlattices grown at Tsub = 500 °C, there is no ZFC/FC discrepancy of superparamagnetism.

The saturation magnetization varies with substrate temperature, shown in Fig. 4(b). As is illustrated in Fig. 7(a), Msat is larger for samples grown at low Tsub. The ferromagnetic hysteresis loops saturate at Msat∼ 0.1 μb/Fe3+ at Tsub = 700 °C, ∼0.2 μb/Fe3+ at Tsub = 600 °C, and ∼0.4 μb/Fe3+ at Tsub = 500 °C [Fig. 4(c)]. The squareness of the ferromagnetic hysteresis also increases at low Tsub. The coercive field (Hc) at 300 K varies between ∼54 and 196 Oe and increases with Tsub, from ∼54 Oe at Tsub = 500 °C to ∼196 Oe at Tsub = 700 °C. The saturation magnetization is also correlated with the RMS surface roughness measurements [Fig. 7(b)]. Superlattices with rougher interfaces have larger values of Msat. Evidently, the ferromagnetic phase increases as the interfaces get rougher. This correlation supports the interfacial magnetism hypothesis for the ferromagnetic response of these superlattices. As the growth surface becomes rougher, and the interface ferromagnetic region increases, the saturation magnetization increases with a larger fraction of Fe in a ferromagnetic state. Conversely, at higher substrate temperatures where the interfaces are smoother, the saturation magnetization decreases.

Additionally, these trends in the ferromagnetic response are consistent with the observations of superparamagnetism as evidenced by the ZFC/FC data described above. Samples with smoother interfaces grown at high temperature exhibit a clear superparamagnetic component (ZFC/FC divergence), which is consistent with a reduction in thickness of the interface ferromagnetic phase. As the thickness of the phase reduces below the stable domain size, superparamagnetic behavior is expected. Conversely, samples with rougher interfaces that exhibit the strongest ferromagnetic response exhibit no clear superparamagnetic behavior. The overall density of paramagnetic moments, as obtained through the Brillouin fits, exhibits a very small decrease with roughness, considering the error bars [Fig. 7(c)]. Although this trend is very weak, it does correlate well with the trends in the superparamagnetic response and the ferromagnetic response, such that rougher interfaces appear to lead not only to a smaller superparamagnetic component but also to a smaller number of paramagnetic moments and a larger ferromagnetic component (high Msat).

The combined structural and magnetic characterization study described above is consistent with the ferromagnetic component in the magnetization of these superlattices emerging from the interface layer between μ-Fe2O3 and β-Ga2O3 and rules out the presence of secondary phases such as Fe clusters. This supports the previous DFT calculations, where ferromagnetism emerges at the interface region of the Ga and Fe phases of the monoclinic μ-Fe2O3/β-Ga2O3 superlattice structures.1 However, in addition to experimentally validating the interface ferromagnetism prediction, we observe the emergence of superparamagnetism in samples with the smoothest interfaces. This suggests that as the ferromagnetic interface layer becomes too thin (as expected for smooth interfaces), the zero-field magnetization becomes unstable against rotations, i.e., superparamagnetism. Further work is needed to examine the size and geometry of these superparamagnetic clusters.

Superlattices composed of ferromagnetic μ-Fe2O3 and diamagnetic β-Ga2O3 on (010) oriented β-Ga2O3 substrates were grown using PAMBE. The interface quality was varied by forming identical superlattice structures at various growth conditions (Tsub = 500–700 °C, ΦGa = 2.5–6 × 10−8 Torr). STEM showed that the μ-Fe2O3 and β-Ga2O3 layers are lattice matched across the entirety of the superlattice. AFM imaging shows that the growth morphology becomes rougher as the substrate temperature and Ga fluxes are decreased. The AFM surface roughness is ∼2.1× larger for the superlattices grown at Tsub = 500 °C than at Tsub = 700 °C. The XRD linewidths are similarly ∼2× larger for the superlattices grown at ΦGa = 2.5 × 10−8 Torr than at ΦGa = 6 × 10−8 Torr. Magnetization data also confirm that the ferromagnetic response changes with the properties of the interfaces. The ferromagnetic component of the magnetization increases with interface roughness in agreement with the interface origin of the ferromagnetic phase in μ-Fe2O3/β-Ga2O3 superlattices. Additionally, samples with a smoother interface exhibiting a smaller ferromagnetic component show a superparamagnetic response and a higher density of paramagnetic moments.

This work was supported by the Center for Emergent Materials at The Ohio State University, an NSF MRSEC (No. DMR-1420451), and by the Army Research Office MURI (No. W911NF-14-1-0016). Electron microscopy was performed at the Center for Electron Microscopy and Analysis (CEMAS) at The Ohio State University.

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