wafer-scale, with good crystallinity and with contamination levels compatible with large-scale back-end-of-line (BEOL) integration. At present, chemical vapor deposition (CVD) on catalytic copper (Cu) substrates is widely recognized as the most promising route to obtain scalable monolayer graphene for electronic and optoelectronic applications. [1][2][3][4] However, significant hurdles are limiting the actual integration of CVD graphene grown on Cu for most applications. In the first instance, the unavoidable transfer process over wafer-scale is rather cumbersome and introduces contamination, unintentional doping, and mechanical stress, [5][6][7] which adversely impact the physical integrity and electrical performance [8] of the graphene layer. The significant challenge involved in carrying out this seemingly straightforward task is reflected by the vast literature on large-scale transfer processes. Second, metallic contamination levels in transferred CVD graphene grown on Cu are typically well-above the specifications requested for BEOL integration. [6] Clearly, asThe adoption of graphene in electronics, optoelectronics, and photonics is hindered by the difficulty in obtaining high-quality material on technologically relevant substrates, over wafer-scale sizes, and with metal contamination levels compatible with industrial requirements. To date, the direct growth of graphene on insulating substrates has proved to be challenging, usually requiring metal-catalysts or yielding defective graphene. In this work, a metal-free approach implemented in commercially available reactors to obtain high-quality monolayer graphene on c-plane sapphire substrates via chemical vapor deposition is demonstrated. Low energy electron diffraction, low energy electron microscopy, and scanning tunneling microscopy measurements identify the Al-rich reconstruction9° of sapphire to be crucial for obtaining epitaxial graphene. Raman spectroscopy and electrical transport measurements reveal high-quality graphene with mobilities consistently above 2000 cm 2 V −1 s −1 . The process is scaled up to 4 and 6 in. wafers sizes and metal contamination levels are retrieved to be within the limits for back-end-ofline integration. The growth process introduced here establishes a method for the synthesis of wafer-scale graphene films on a technologically viable basis.

Nonpolar (112¯0) a-plane GaN films have been grown by metal-organic vapor deposition on r-plane (11¯02) sapphire. Lateral growth is favored using a low V:III ratio resulting in films with a smooth surface, while pitted films are grown at a high V:III ratio indicating preferential on-axis growth. High-resolution x-ray diffraction analysis of both film types showed a strong anisotropy in the peak width of the symmetric omega rocking curve with respect to the in-plane orientation, phi. In-plane isotropic behavior of crystallinity with overall reduced omega full width at half maximum values was achieved when the growth was initiated at a high V:III ratio before reducing the V:III ratio for film coalescence. An improvement of crystal quality through initial surface roughening was equally realized by the incorporation of partial-coverage SiNx interlayers.

We have studied a series of GaN films grown with a range of dislocation densities by atomic force microscopy (AFM), transmission electron microscopy (TEM) and high resolution x-ray diffraction (HRXRD). The (002), (004), (006), (105), (204), (302), (100), (110), (200) and (300) reflections were measured as reciprocal space maps (RSMs) or scans in ω and ω/2θ. The latter 4 in-plane reflections were measured using a low, or glancing, incident angle with respect to the film surface. We have used a variety of different methods to try and obtain reliable measurements for mosaic tilt, twist, crystallite size and microstrain both in- and out-of plane. From (hk0) data in-plane twist angles were measured ranging from 0.37° to 0.078° and in-plane microstrains from 3.5 × 10−4 to 1.8 × 10−4. The improvements in the quality of the GaN layers relate to the increased island coalescence time, which reduces in particular the number of edge-type threading dislocations. The first three samples had a much larger tilt ∼0.09° than the last three ∼0.04°. However, the latter samples were bowed, so results from a single measurement on the (002) peak are too large. Beam restriction on the (002) or an extrapolation from several (00l) reflections gives more reliable results. The values obtained for in-plane crystallite size are in general variable or unreliable. For some samples the sizes are considered to be too large to be accessible by XRD; in most cases the peak broadening is dominated by tilt or twist or microstrain and the results are sensitive to assumptions about the peak shape. For the samples with smaller measurable crystallite sizes, the (hk0) peaks were too weak to measure reliably. The cell parameters showed more compressive strain with fewer dislocations. The trends observed by HRXRD are consistent with AFM and TEM results.

There is a great deal of interest in the underlying causes of efficiency droop in InGaN/GaN quantum well light emitting diodes, with several physical mechanisms being put forward to explain the phenomenon. In this paper we report on the observation of a reduction in the localization induced S-shape temperature dependence of the peak photoluminescence energy with increasing excitation power density. This S-shape dependence is a key fingerprint of carrier localization. Over the range of excitation power density where the depth of the S shape is reduced, we also observe a reduction in the integrated photoluminescence intensity per unit excitation power, i.e., efficiency droop. Hence, the onset of efficiency droop occurs at the same carrier density as the onset of carrier delocalization. We correlate these experimental results with the predictions of a theoretical model of the effects of carrier localization due to local variations in the concentration of the randomly distributed In atoms on the optical properties of InGaN/GaN quantum wells. On the basis of this comparison of theory with experiment we attribute the reduction in the S-shape temperature dependence to the saturation of the available localized states. We propose that this saturation of the localized states is a contributory factor to efficiency droop whereby nonlocalized carriers recombine non-radiatively. V

Two-dimensional (2D) crystals have renewed opportunities in design and assembly of arti cial lattices without the constraints of epitaxy. However, the lack of thickness control in exfoliated van der Waals (vdW) layers prevents realization of repeat units with high delity. Recent availability of uniform, waferscale samples permits engineering of both electronic and optical dispersions in stacks of disparate 2D layers with multiple repeating units. We present optical dispersion engineering in a superlattice structure comprised of alternating layers of 2D excitonic chalcogenides and dielectric insulators. By carefully designing the unit cell parameters, we demonstrate > 90 % narrowband absorption in < 4 nm active layer excitonic absorber medium at room temperature, concurrently with enhanced photoluminescence in cm 2 samples. These superlattices show evidence of strong light-matter coupling and exciton-polariton formation with geometry-tunable coupling constants. Our results demonstrate proof of concept structures with engineered optical properties and pave the way for a broad class of scalable, designer optical metamaterials from atomically-thin layers.