Magnetic exchange field in magnetic multilayers can potentially reach tens or even hundreds of Tesla. 6 The single-atomic-layer (2D) materials, such as graphene, mono-layer WS 2 etc., is expected to experience the strongest MEF in heterostructures with magnetic insulators due to the short-range nature of magnetic exchange coupling. 4 2D material/magnetic insulator heterostructures enable local spin modulation by magnetic gates, 4,5,7 and the realization of efficient spin generation for spintronic applications. 8,9 As a proof of concept, here we demonstrate substantial MEF and spin polarization in CVD graphene/EuS heterostructures. We have chosen EuS as a model magnetic insulator because of its wide band-gap (1.65 eV), large exchange coupling J~10 meV, and large magnetic moment per Eu ion ௭~7 , 10 yielding large estimated exchange splitting ௭ in graphene. 4,5 EuS has also been shown to spin-polarize quasiparticles in materials including superconductors and topological insulators. 6,11 The strength of the MEF depends critically on the interface and EuS quality, 12,13 which we optimize with an in-situ cleaning and synthesis process (Methods and Fig. 1a). In contrast to other means, such as defect-or adatom-induced spin polarization, 14,15 depositing insulating EuS well preserves graphene's chemical bonding, confirmed by Raman spectroscopy (Fig. 1b) (Fig. S5-1), indicative of high graphene quality and well-preserved Dirac band structure.We utilize Zeeman spin-Hall effect (ZSHE) to probe the MEF in graphene which splits the Dirac cone via Zeeman effect and generates electron-and hole-like carriers with opposite 4 spins near the Dirac point ( Fig. 2a right panel). 8,9 Under a Lorentz force, these electrons and holes propagate in opposite directions, giving rise to a pure spin current and non-local voltage ( Fig. 2a left panel). We measure the non-local resistance of ZSHE using the device configuration in Fig. 2a where ௫ is the MEF. We further define the parameter :where ௭ denotes the Zeeman energy at the reference field . Given , deriving of graphene/AlO x is straightforward because ௭ is solely determined by . The inset of Fig. 3(b) shows the calculated using T, a proper reference field as we will explain below.To derive of graphene/EuS, we note that according to the theory of ZSHE, 9,17 depends on sample mobility, while other sample-dependents terms (including spin relaxation length, density of thermally activated carriers and Fermi velocity) cancel out (see S3 in SI). The mobility difference between our graphene/EuS and graphene/AlO x samples is~25% (see S1 in SI), which would only yield a~10% correction to (see S3 in SI). Since~10% difference is 6 small, for an order-of-magnitude estimate of the MEF, we adopt the value of graphene/AlO x for graphene/EuS as an approximation. We then evaluate E Z in graphene/EuS usingTo obtain the lower bound of , we approximate , ignoring the ௫contribution. This constrains us to use a small such that ௫ is small. Meanwhile, should be high enough to ensure that , is much large...
Transition metal dichalcogenides (TMDs) have emerged as promising materials to complement graphene for advanced optoelectronics. However, irreversible degradation of chemical vapor deposition-grown monolayer TMDs via oxidation under ambient conditions limits applications of TMD-based devices. Here, the growth of oxidation-resistant tungsten disulfide (WS ) monolayers on graphene is demonstrated, and the mechanism of oxidation of WS on SiO , graphene/SiO , and on graphene suspended in air is elucidated. While WS on a SiO substrate begins oxidation within weeks, epitaxially grown WS on suspended graphene does not show any sign of oxidation, attributed to the screening effect of surface electric field caused by the substrate. The control of a local oxidation of WS on a SiO substrate by a local electric field created using an atomic force microscope tip is also demonstrated.
We demonstrate an electrochemical method-which we term oxidative decoupling transfer (ODT)for transferring chemical vapor deposited graphene from physically deposited copper catalyst layers. This copper oxidation-based transfer technique is generally applicable to copper surfaces, and is particularly suitable where the copper is adhered to a substrate such as oxidized silicon. Graphene devices produced via this technique demonstrate 30% higher mobility than similar devices produced by standard catalyst etching techniques. The transferred graphene films cover more than 94% of target substrates-up to 100 mm diameter films are demonstrated here-and exhibit a low Raman D:G peak ratio and a homogenous and continuous distribution of sheet conductance mapped by THz time-domain spectroscopy. By applying a fixed potential of-0.4V vs. an Ag/AgCl reference electrode-significantly below the threshold for hydrogen production by electrolysis of water-we avoid the formation of hydrogen bubbles at the graphene-copper interface, preventing delamination of thin sputtered catalyst layers from their supporting substrates. We demonstrate the reuse of the same growth substrate for five growth and transfer cycles and prove that this number is limited by the evaporation of Cu during growth of graphene. This technique therefore enables the repeated use of the highest crystallinity and purity substrates without undue increase in cost.
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In thermionic energy converters, the absolute efficiency can be increased up to 40% if space-charge losses are eliminated by using a sub-10-µm gap between the electrodes. One practical way to achieve such small gaps over large device areas is to use a stiff and thermally insulating spacer between the two electrodes. We report on the design, fabrication and characterization of thin-film alumina-based spacers that provided robust 3–8 μm gaps between planar substrates and had effective thermal conductivities less than those of aerogels. The spacers were fabricated on silicon molds and, after release, could be manually transferred onto any substrate. In large-scale compression testing, they sustained compressive stresses of 0.4–4 MPa without fracture. Experimentally, the thermal conductance was 10–30 mWcm−2K−1 and, surprisingly, independent of film thickness (100–800 nm) and spacer height. To explain this independence, we developed a model that includes the pressure-dependent conductance of locally distributed asperities and sparse contact points throughout the spacer structure, indicating that only 0.1–0.5% of the spacer-electrode interface was conducting heat. Our spacers show remarkable functionality over multiple length scales, providing insulating micrometer gaps over centimeter areas using nanoscale films. These innovations can be applied to other technologies requiring high thermal resistance in small spaces, such as thermophotovoltaic converters, insulation for spacecraft and cryogenic devices.
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