Monolithic integration of optoelectronics with electronics is a much-desired functionality. Here, we demonstrate that it is possible to realize low-loss Ge quantum-well photonic interconnects on Si wafers. We show that Ge-rich Si1–xGex virtual substrates can act as a passive, high-quality optical waveguide on which low-temperature, epitaxial growth of Ge quantum-well devices can be realized. As a proof of concept, the photonic integration of a passive Si0.16Ge0.84 waveguide and two Ge/SiGe multi-quantum-well active devices, an optical modulator and a photodetector was realized to form a photonic interconnect using a single epitaxial growth step. This demonstration confirms that Ge quantum-well interconnects are feasible for low-voltage, broadband optical links integrated on Si chips. Our approach can be extended to any kind of Ge-based optoelectronic device working within telecommunication wavelengths as long as a suitable Ge concentration is selected for the Ge-rich virtual substrate.
Graphene can be transformed from a semimetal into a semiconductor if it is confined into nanoribbons narrower than 10nm with controlled crystallographic orientation and well-defined armchair edges. However, the scalable synthesis of nanoribbons with this precision directly on insulating or semiconducting substrates has not been possible. Here we demonstrate the synthesis of graphene nanoribbons on Ge(001) via chemical vapour deposition. The nanoribbons are self-aligning 3° from the Ge110 directions, are self-defining with predominantly smooth armchair edges, and have tunable width to <10nm and aspect ratio to >70. In order to realize highly anisotropic ribbons, it is critical to operate in a regime in which the growth rate in the width direction is especially slow, <5nmh−1. This directional and anisotropic growth enables nanoribbon fabrication directly on conventional semiconductor wafer platforms and, therefore, promises to allow the integration of nanoribbons into future hybrid integrated circuits.
Graphene nanoribbons are excellent charge1 and thermal2 conductors that can exhibit high current-carrying capacity3 and novel magnetic and spin-polarized edge states4, 5, 6 depending on their crystallographic orientation, edge structure and width. Unlike continuous two-dimensional graphene, which is semimetallic, one-dimensional (1D) graphene nanoribbons can be semiconducting, allowing for substantial modulation of their conductance and enabling their application in semiconductor logic, high-frequency communication devices, optoelectronics, photonics and sensors in which a bandgap is needed to achieve high performance. Their bandgap roughly varies inversely with ribbon width and the largest bandgaps are expected for ribbons with armchair edge orientation7.
While achieving a ribbon width of <10nm is necessary to induce a technologically relevant bandgap that is substantially greater than kBT of 25meV at room temperature, 7, sub-10nm resolution is beyond the limits of conventional optical and electron-beam lithography. Moreover, top-down lithographic techniques in which ribbons are etched from continuous graphene sheets result in nanostructures with relatively rough, defective edges, which lead to Coulomb blockade8 and localized electronic9, 10 and phonon states11 and, consequently, degrade the high charge carrier mobility12, 13, 14 and thermal conductivity2 of graphene.
These deficiencies, in part, can be overcome, via bottom-up organic synthesis on metal surfaces15,16, 17 and in solution18, 19 as well as by unzipping graphite20 and carbon nanotubes21 in solution, to yield ribbons with sub-10nm width and smooth edges. However, surface-assisted organic synthesis yields short ribbons (~20nm) and unzipping graphite and carbon nanotubes does not offer control over the ribbon crystallographic orientation. Moreover, the controlled placement and alignment of ribbons onto substrates from solution has proven to be difficult.
Scalable nanoribbon fabrication has also been reported via epitaxial growth on templated SiC nanofacets22 and by chemical vapour deposition (CVD) on surface features, such as steps23, twins24 and trenches25, as well as on patterned catalysts in which growth is confined to predetermined areas that define the ribbon dimensions26, 27, 28, 29. However, with these approaches, ribbons with sub-10nm width have not been demonstrated and the catalyst template determines the ribbon edge structure rather than a more precise self-defining growth mechanism.
Here we show that the CVD of graphene on Ge(001) can be controlled to yield oriented nanoribbons with sub-10nm width and smooth armchair edges. Previous work on integrating graphene with Ge has primarily focused on large-area graphene monolayers. For example, continuous graphene films have been transferred onto Ge from other substrates30, 31. Furthermore, continuous graphene monolayers have been grown via CVD directly on Ge(001) by Wang et al.32 and Ge(110) and Ge(111) by Lee et al.33. However, in these previous studies, nanoribbons were not observed in partial growth experiments. In this work, we report that high aspect ratio nanoribbons can be directly grown on the Ge(001) facet by tailoring the CVD conditions to maximize the anisotropy of crystal growth. It is critical to operate in a regime in which the growth rate is especially slow, <5nmh−1 in the width direction. Nanoribbons are grown at atmospheric pressure using various growth temperatures (860<T<935°C), CH4 mole fractions (3.7 × 10−3< <1.6 × 10−2), and H2 mole fractions (0.17< <0.33), as summarized inSupplementary Table 1. By tuning T, , and the growth time (t), the growth anisotropy is tailored to yield ribbons from the bottom-up with controlled width (w), length (l) and aspect ratio. Using conditions in which the anisotropy is maximized, isolated ribbons are prevalent on the Ge(001) surface even after t>18h. In contrast, Wang et al.32 used a relatively fast growth rate in which continuous graphene films were synthesized on Ge(001) in t<100min.
Several general observations are made regarding graphene growth on Ge(001) by analysing representative scanning electron microscopy (SEM), atomic force microscopy (AFM) and scanning tunnelling microscopy (STM) images in Fig. 1a–c. Following nucleation, graphene crystals evolve anisotropically, resulting in nanoribbons with high aspect ratio and smooth, straight edges. Raman spectroscopy indicates that the 2D:G ratio and the 2D peak full-width-at-half-maximum are 6.0 and 28cm−1, respectively, confirming that the nanoribbons are monolayer graphene34 (Supplementary Fig. 1). The ribbons preferentially orient closely along the 110 directions of the Ge(001) template, resulting in two ribbon orientations that are approximately perpendicularly aligned. These two orientations exist with equal probability. For ribbons with w<10nm, the short ribbon edges form 60, 90 or 120° angles with the long ribbon edges. However, for wider ribbons, only angles of 60 and 120° are observed, indicating that all edges are oriented along equivalent crystallographic directions of graphene. While as high as 90% of the graphene crystals that nucleate evolve as ribbons, more compact graphene crystals with lower aspect ratio and edges that are not aligned along Ge110 directions are also observed. Interestingly, the Ge underneath the ribbons is nanofaceted, which is studied further below.
Growth kinetics and evolution
We quantify the growth kinetics to gain insight into the processes that determine the ribbon size and aspect ratio using constant T of 910°C, of 0.0092 and of 0.33. Both w and l increase witht, along with the ribbon-to-ribbon variation in w and l, which is quantified by the range of the box and whiskers in Fig. 1d,e. The mean growth rates in the w and l directions, Rw and Rl, respectively, are compared in the insets of Fig. 1d,e (where w and l increase on average at twice Rw and Rl). Initially, Rl is 90nmh−1 whereas Rw is only 5nmh−1, giving rise to the anisotropic ribbon evolution. While Rl is relatively constant with time, Rw increases to >10nmh−1 after several hours. Accordingly, the mean aspect ratio decreases from 20 to 10 with increasing t (Fig. 1f). Stopping growth after t of 1h results in ribbons with average w of 9.8nm (Fig. 1g), which is below the resolution of optical and typical electron-beam lithography, and we anticipate that even narrower ware obtained at earlier t. Figure 1c shows an example of a nanoribbon with w of 7nm and l of 160nm.
The effects of , and T on the nanoribbon growth rate and the resulting anisotropy are also quantified. The growth rates increase as increases (Fig. 2a), decreases (Fig. 2b), and T increases (Supplementary Fig. 2). Each of these parameters can be independently tuned to vary Rl over an order of magnitude from 30 to 300nmh−1. The anisotropy varies inversely with Rl, independent of whether or is changed (Fig. 2c). Thus, a critical parameter for realizing high aspect ratio nanoribbons is to operate in a regime in which growth is slow. For example, at low of 0.0066 and high of 0.33, not only is a slow Rl of 40nmh−1 achieved, but also a much slower Rw of 1.4nmh−1, yielding ribbons with an aspect ratio of 30, on average, and as high as 70 (Supplementary Fig. 3). Minimizing Rw also makes it possible to tailor w with high precision. Empirical rate laws are identified (insets of Fig. 2a,b), indicating that Rl scales as and . Furthermore, an Arrhenius temperature dependence is observed with an activation energy of 7.2±0.4eV (Supplementary Fig. 2). For comparison, the activation energy for graphene growth on Cu is only 1–3eV (ref. 35). However, graphene growth on Cu at atmospheric pressure yields hexagonal crystals instead of nanoribbons36, highlighting that different mechanisms control growth.
We also find that the 1D nature of growth is insensitive to the bulk Ge dopant concentration (NSb<1.5 × 1018cm−3), Ge surface treatment prior to synthesis (OH, H and Cl functionalization), and annealing time before growth (Supplementary Figs 4–6). Nanoribbons are not observed on Ge(110) nor Ge(111) under any growth condition (Supplementary Fig. 7).
Low-energy electron microscopy and diffraction
Low-energy electron microscopy (LEEM) and diffraction (LEED) are used to determine the orientation of the graphene lattice with respect to the underlying Ge and with respect to the ribbon edges. The overlaid LEED patterns in Fig. 3a taken at 121 and 135eV establish the Ge and directions. The LEED pattern in Fig. 3b obtained at 67eV shows that the graphene lattice primarily exists within two families of crystallographic orientations, denoted by the orange and purple hexagons. Dark-field imaging in Fig. 3c indicates that the purple family of diffraction spots originates from ribbons that are oriented with their long axis approximately parallel to Ge whereas the orange family of spots originates from ribbons that are oriented with their long axis approximately parallel to Ge . Both of these families of ribbons have edges that are macroscopically aligned with the armchair direction of graphene. This armchair edge orientation is unique because graphene crystals grown on Cu and Ni typically have edges that are aligned along the zigzag direction of graphene25, 37.
Within each family, there are two unique graphene orientations that are rotated 2.9±0.4° (~3°) relative to the Ge110 directions, as indicated by the diffraction spots belonging to the orange family that are circled in red and blue in Fig. 3b. These nanoribbon orientations are depicted in Fig. 3d. Dark-field imaging in Fig. 3c indicates that some of the ribbons are single crystalline, corresponding to either the +3° or −3° graphene orientation, whereas others are bi-crystalline, in which the crystal lattice of one half of the ribbon is rotated by 2 × 3°=6° with respect to the other half. This indicates that the nanoribbons nucleate in their centre and then grow in opposite directions along their length. These dynamics potentially explain why the ribbons are elevated in their centre (Fig. 1c); the sublimation of Ge may be locally suppressed under the ribbons as they grow. Interestingly, the lattice of the non-ribbon graphene crystals with lower aspect ratio and edges not aligned on Ge110, like the one observed in Fig. 1a, is rotated with respect to the lattice of the nanoribbons, typically by 15° as observed in Fig. 3b. This difference indicates that the anisotropic nanoribbon growth is driven only when there is a specific relative orientation between the graphene lattice and the Ge(001) surface.
Scanning tunneling microscopy
Ultra-high vacuum STM is performed to substantiate the LEED data and to determine the atomic nature of the graphene nanoribbon edges on Ge(001). The STM image and its corresponding fast Fourier transform (FFT) in Fig. 4a indicate that the ribbon edges are straight and parallel to the armchair direction of graphene with little edge roughness and that the graphene lattice is rotated 3° from the Ge110 directions, consistent with the LEED data. The Ge underneath the nanoribbons retains the common (2 × 1) dimer reconstruction (Fig. 4b–d) even after ambient exposure. Contamination of the bare Ge surface upon exposure to ambient often precludes precise topographical imaging of the nanoribbon edge structure. Using the topographic data alone, we can set an upper limit on the edge roughness. For example, the representative 40nm ribbon segment inFig. 4a has roughness of <0.5nm (two lattice constants of graphene). However, we can learn more about the edge structure from quantum interference patterns caused by intervalley backscattering (Fig. 4e) of charge carriers at the ribbon edges. The ring-like shapes with a R30° unit cell highlighted with the rhombuses in Fig. 4g and the armchair-like patterns with periodicity (λf) of 3.7Å in Fig. 4g are consistent with electron backscattering at armchair edges16, 38 and are clearly revealed by the prominence of the K/K′ points in the FFT in Fig. 4f. The presence of these coherent interference patterns combined with the small line edge roughness indicates that the edges consist primarily of smooth armchair segments. The interference patterns decay into the interior of the ribbons with length scales comparable to graphene on SiC38 and metals39. The atomic structure of the hexagonal graphene lattice is observed past these decay lengths in the interior of the nanoribbons (Fig. 4g).