A Study of GaAs1âxSbx Axial Nanowires Grown on Monolayer

Article Cite This: ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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A Study of GaAs1−xSbx Axial Nanowires Grown on Monolayer Graphene by Ga-Assisted Molecular Beam Epitaxy for Flexible NearInfrared Photodetectors Surya Nalamati,† Manish Sharma,‡ Prithviraj Deshmukh,‡ Jeffrey Kronz,§ Robert Lavelle,§ David Snyder,§ C. Lewis Reynolds, Jr.,∥ Yang Liu,∥ and Shanthi Iyer*,†,‡

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Department of Electrical and Computer Engineering and ‡Nanoengineering, Joint School of Nanoscience and Nanoengineering, North Carolina A&T State University, Greensboro, North Carolina 27411, United States § Penn State Applied Research Laboratory, The Pennsylvania State University, University Park, Pennsylvania 16802, United States ∥ Department of Materials Science and Engineering, North Carolina State University, Raleigh, North Carolina 27695, United States S Supporting Information *

ABSTRACT: We report the successful growth of high-quality GaAs1−xSbx nanowires on monolayer graphene/SiO2/p-Si (111) using molecular beam epitaxy (MBE) for the application of a flexible near-infrared photodetector. A systematic and detailed study of NW growth parameters, namely, growth temperature, V/III beam equivalent pressure (BEP) ratio, and Ga shutter opening duration, has been carried out. Growth of vertical ⟨111⟩ oriented nanowires on graphene with 4 K photoluminescence emission in the range 1.24−1.38 eV has been achieved. The presence of a weak D mode in Raman spectra of NWs grown on graphene suggests that NW growth did not alter the intrinsic properties of the monolayer graphene. High-resolution transmission electron microscopy and a selective area diffraction pattern confirmed the zinc-blende crystal structure of the NWs. This study suggests that Sb as a surfactant plays a critical role in the surface engineering of the substrate, leading to the superior optical quality of NWs exhibiting a higher 4 K photoluminescence intensity and lower full width at half maxima (fwhm) with significant improvement in optical responsivity compared to NWs grown on Si substrate of similar Sb composition. KEYWORDS: nanowire, graphene, GaAsSb, molecular beam epitaxy, current sensing atomic force microscopy, photodetector

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INTRODUCTION

attributes of both the substrate and the NWs. Among the 2D substrate materials, graphene has stimulated extensive research over the past decade due to its exceptional electrical, mechanical, and optical properties while being cost-effective and relatively easy to produce.17−19 Graphene is also a candidate for integration into next-generation electronics to continue scaling the device density following Moore’s law. Additionally, it is an ideal platform for flexible devices and enables the exploration of novel designs,20 which makes use of optical transparency on both sides of the devices to meet future energy harvesting demands.21−23 One desirable attribute of III−V material growth on graphene is the hexagonal symmetry of the (111) plane which is identical to the alignment of carbon atoms in the structure of graphene. Therefore, this substrate is well-suited for the preferred NW growth in the ⟨111⟩ direction. Lastly, there are three preferred sites of adsorption on graphene that are identified as H, B, and T, as opposed to no such preferred

Semiconductor nanowires (NWs) have stimulated great interest due to their large aspect ratio, which enables relaxation of lattice mismatch constraints and enhanced light trapping features. This offers unique opportunities for integrating nanophotonic devices, primarily based on III−V compound semiconductors with electronic devices on a Si platform as well as other nanoscopic and microscopic systems.1 Among III−V NWs, GaAsSb-based NWs have attracted considerable attention due to the wide bandgap tunability from 870 nm (GaAs) to 1700 nm (GaSb),2−4 which encompasses the wavelength range of interest to optical communications with a high absorption coefficient and superior carrier mobility characteristics.2,5,6 These features are attractive for a wide range of applications at the nanoscale as in nanophotonic integrated circuits,4,7 photodetectors,4 lasers,8,9 optical telecommunications,10 and quantum information science.7 GaAsSb NWs on (111) Si have been widely investigated.1,2,6,11−16 The ability to tolerate larger mismatch by the NWs offers an opportunity to investigate the growth of these NW materials on other types of substrates, such as 2D materials that enables taking advantage of the distinct © XXXX American Chemical Society

Received: May 11, 2019 Accepted: May 31, 2019 Published: May 31, 2019 A

DOI: 10.1021/acsanm.9b00893 ACS Appl. Nano Mater. XXXX, XXX, XXX−XXX

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ACS Applied Nano Materials sites present on Si.24,25 The former two sites are more favorable for III−V adatoms, and the preferred site among these two depends on the type of adatoms.26 Thus, different possible combinations of lattice parameters exist depending on the type of atoms and the arrangement based on the NW material composition.17,24 Although growing NWs on graphene seems very attractive and even promising, it can also be challenging. For example, a major challenge is graphene’s low surface energy,27 which is 2 orders of magnitude lower than conventional substrates such as Si and GaAs28 due to its sp2 hybridization. Because the conventional method of growing NWs, the vapor−liquid−solid (VLS) mechanism, depends on surface energy,29 graphene’s low surface energy leads to dewetting of the droplet on the substrate due to the lack of any significant interaction between the droplet and the substrate.30 The formation of the droplet thus formed with a large wetting angle is not conducive for subsequent NW growth.31 Hence, there are limited reports on the growth of semiconductor NWs on graphene.20,21,24,32−34 The growth of GaAs NWs on multilayer graphene,24 GaN NWs on single and multilayer graphene,35 InAs NWs on a monolayer of graphene,17 InAsSb on graphite,36 and ZnO NWs on graphene foam33 have all been reported. As is evident, growing NWs on graphene is still in the nascent stage, which is why there has been little progress report on Sb-based III−V semiconductor NWs grown on graphene. In this work, we report on the growth of GaAs1−xSbx axial NWs on monolayer graphene by Ga-assisted molecular beam epitaxy (MBE) for the application of a flexible near-infrared (NIR) photodetector. A detailed and systematic study has been conducted to correlate growth parameter variations, namely, growth temperature, V/III beam equivalent pressure (BEP) ratio, and Ga shutter opening duration, to improvement in NW quality, vertical alignment, and Sb incorporation in the NWs. This series of investigations also provided insight into the possible role of the Sb as a surfactant on the surface engineering of the graphene. The results reveal that Sb has a substantial impact on the engineering of the graphene surface, resulting in possible diameter modulation of the NWs, more significant influence of V/III BEP ratio on the Sb incorporation in the NWs, and improved quality of GaAsSb NWs resulting in enhanced optical responsivity in comparison to those of similar composition grown on Si. These results are significant since they may be broadly applicable to growing other high-quality semiconductor NWs with one of their constituent elements being a surfactant on graphene. This may open other avenues for designing high-performing optoelectronic devices at the nanoscale.

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Afterward, the Cu was etched by using CE-100 FeCl3 Cu etchant, followed by DI water rinse, 10% HCl rinse, and a final DI water rinse (15 min each). Pristine single-layer graphene (SLG) of 100 ± 0.3 mm diameter was then transferred by using a wet transfer method onto 15 nm thick SiO2, which was grown by a thermal oxidation process on a p-Si (111) substrate. The root-mean-square (rms) surface roughness of SiO2 was ∼1 nm. The transferred sample is baked at 50 °C for 4 h and then 10 min at 150 °C to remove any moisture between the film and substrate. The PMMA is then stripped with the acetone and IPA, followed by another 30 min bake at 50 °C to dry. Just before loading into the MBE entry chamber, the transferred SLG on SiO2/p-Si (111) is rinsed with acetone and dried with N2. It is further baked under high vacuum in the intro chamber of the MBE system at 200 °C for 8 h. The growth of GaAs1−xSbx axial NWs13,14,12,37 was carried out by using As4 and Sb2 as the group V constituent sources. The growth was initiated at a lower growth temperature in the range 520−550 °C by opening the Ga shutter, followed by the opening of the As and Sb shutters. The growth temperature was immediately increased to a higher temperature range of 580−610 °C for a growth duration of 5 min. For optimized growth conditions, Ga, As, and Sb BEP of 1 × 10−7, 1.8 × 10−6, and 2 × 10−7 Torr, respectively, for a V/III BEP ratio of 20 were used. The temperature of the Ga effusion cell was preset to yield a nominal planar GaAs growth rate of 0.5 ML/s for Ga BEP of 1 × 10−7 Torr with the As cracker valve cell position set for the desired flux. For other BEP ratios, the group V flux was varied by keeping the Ga BEP invariant. A schematic of our overall graphene and nanowire growth sequence is shown in Figure 1.

Figure 1. Schematic illustration of the growth process of GaAsSb NWs on monolayer graphene. Scanning electron microscopy (SEM) was performed using a Carl Zeiss Auriga-BU FIB field emission scanning electron microscope (FESEM). X-ray diffraction (XRD) was completed by using a Bruker D8 Discover instrument with a DaVinci diffractometer in the standard Bragg−Brentano parafocusing configuration. Scanning transmission electron microscopy (STEM) and energy dispersive X-ray spectroscopy (EDS) characterizations were performed on an aberrationcorrected (probe) FEI Titan G2 system. Selected area diffraction (SAED) and high-resolution transmission electron microscopy (HRTEM) were performed on a JEOL 2010F microscope operated at 200 kV. Optical measurements included μ-photoluminescence (μPL) using a 633 nm He−Ne laser as the excitation source with a 0.32 m double grating monochromator for wavelength dispersion and an InGaAs detector for detection using a conventional lock-in amplifier. A closed-cycle optical cryostat from Montana Cryostation with the sample chamber interfaced with a fiber-coupled confocal microscope was used for μ-PL measurements at 4 K. The sample was dissolved in isopropanol, sonicated for 100 s, and further dispersed on a Rh flashed copper grid of 150 mesh and 3.05 mm o.d. These NW samples were then used for single nanowire (SNW) PL measurements. Raman spectroscopy was performed in a Horiba Jobin Yvon ARAMIS Raman microscope with a He−Ne laser (633 nm) excitation source at room

EXPERIMENTAL DETAILS

Graphene was grown via a catalytic CVD reaction on 99.8% Cu foil. The surface of the Cu before growth is cleaned with semiconductor grade acetone, isopropyl alcohol (IPA), deionized water (DI), and acetic acid, followed by an H2 etch at 1000 °C to ensure a clean surface prior to introducing CH4. Growth occurred in a horizontal hot-wall furnace at a pressure of 1 Torr and heated to 1050 °C. This temperature and pressure were maintained during the hydrogen anneal before growth followed by the introduction of CH4 as the carbon source to initiate growth. The Cu foil cooled naturally in an Ar/H2 environment after the growth process. The graphene is vacuum sealed immediately after growth and then coated with poly(methyl methacrylate) (PMMA) in a cleanroom environment to reduce the risk of particle contamination prior to transfer. B


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ACS Applied Nano Materials

Figure 2. Tapping mode AFM surface topography image of monolayer graphene on SiO2/p-Si (111) substrate: (a) topography and (b) 3D view of phase image. temperature. Surface topography measurements were performed by tapping mode in an Agilent LS 5600 atomic force microscope (AFM), and Si probes were used at a resonant frequency of 190 kHz with an image scanning speed of 0.2 lines/s. Current sensing atomic force microscopy (CSAFM) was performed for obtaining I−V characteristics from a single NW on the grown substrate by using the aforementioned AFM. A Pt/Ir-coated Si cantilever of radius ∼20 nm and spring constant of 0.2 N/m was used. The AFM probe was grounded, and the voltage was supplied to the substrate. A tungsten carbide electrode was clamped on the surface of the substrate before imaging. Initially, Z-image and conductive mapping over an area of 20 μm × 20 μm ensemble NWs was completed. By choosing the appropriate set point, and other scanning parameters, we identified NWs by circles in the Z-image. I−V spectroscopy was performed by placing the AFM tip on the highlighted circles. Two helium−neon lasers with wavelengths of 633 and 673 nm and power densities of ∼5 W/cm2 and ∼0.1 kW/cm2, respectively, were used for optical illumination. Finite element modeling (FEM) of the NW was then performed using Comsol Multiphysics software with Poisson’s and ambipolar transport equations to obtain the electric potential and the carrier concentrations at each simulation point to extract the electrical transport parameters from the best fit of the simulated data to the experimental I−V characteristics. The nanowire was modeled in 2D axisymmetric dimension as a 2D rectangle with cylindrical symmetry. Dimensions of the NWs were taken from SEM measurements. Additional details on the modeling are provided in ref 38.

Table 1. Growth Parameters and Associated Nomenclature of the GaAs1−xSbx NWs Grown on Graphene/SiO2/p-Si (111) Substrate sample

Tg1 (°C)

Tg2 (°C)

V/III BEP ratio

Ga shutter opening time (s)

A B(20) C D E(18) F(15) G′(20) H′(18) I′(15) J′(25)

520 540 550 540 540 540 540 540 540 540

580 580 580 610 580 580 580 580 580 580

20 20 20 20 18 15 20 18 15 25

10 10 10 10 10 10 0 0 0 0

In the two step-growth approach, initial substrate temperatures (Tg1) from 520 to 550 °C, representing the initiation of GaAsSb NWs’ growth, were examined. The temperature was then gradually increased during growth to a higher temperature (Tg2) of 580 °C. For these growths, a V/III BEP ratio of 20 with a Ga shutter opening duration of 10 s before growth was used. The NWs grown at Tg1 (520 °C) exhibited nonuniform length, as shown in Figure 3a, whereas an increase in Tg1 to 540 °C yielded NWs of well-defined hexagonal symmetry with an increase in both axial and radial growth rates and NW density, as shown in Figure 3b. A further increase in Tg1 to 550 °C showed an adverse effect on the growth rate and NW density with an increase in “crawling” NWs (see Figure 3c). This is attributed to an increase in the Ga adatom diffusion length with growth temperature and a higher probability of the Ga adatom occupying the unfavorable B or T sites.24,26 Furthermore, parasitic growth was found to be enhanced with an increase in Tg1, which is attributed to nonwetting of the Ga droplet. With the optimized initial growth temperature Tg1 of 540 °C, the influence of Tg2 (580 and 610 °C) on growth rate and NW density was investigated. An increase in Tg2 to 580 °C led to well-faceted NWs with some parasitic growth (Figure 3b). A further increase in Tg2 to 610 °C yields short NWs with very low density and large Ga droplets on the surface with no other visible parasitic growth as displayed in Figure 3d. All these observations are summarized graphically in Figure 3e,f.

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RESULTS AND DISCUSSION An AFM tapping mode topography image of the wrinkle-free surface of the monolayer graphene transferred on the SiO2/Si, as shown in the Figure 2a, exhibits a rms roughness of ∼0.4 nm and a step height of ∼3 nm to the graphene surface. The graphene is polycrystalline, and the average grain size is on the order of 1−2 μm. The size of PMMA residues on graphene is typically ∼1−2 μm, as shown in Figure 2b. These may limit the adatom diffusion but are not expected to alter significantly the NW growth and interface properties.39 The techniques developed for complete removal of PMMA residues on transferred graphene have their own limitations.40 However, the low intensity of the Raman D-peak shown later indicates the absence of any unusual high density of defects in graphene. In the following section, the effects of a two-step growth temperature approach, variation of V/III BEP ratio, and Ga shutter opening duration on the NW growths are presented. Table 1 lists the nomenclature of the samples and associated growth conditions. C


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Figure 3. Variations of surface morphology of GaAsSb axial NWs for different Tg1 with Tg2 at 580 °C: (a) 520 °C (sample A), (b) 540 °C (sample B′(20)), and (c) 580 °C (sample C) and for different Tg2 with Tg1 at 540 °C: (d) 610 °C (sample D). The dependencies of NW lengths, diameters, and density on Tg1 and Tg2 are summarized in graphs (e) and (f), respectively. The error bars represent the extrema of the measurements taken from 20 vertical NWs.

Figure 4. SEM images of GaAsSb axial NWs with variation in V/III BEP ratio of (a) 15 (sample F(15)), (b) 18 (sample E(18)), and (c) 20 (sample B(15)) for a Ga shutter opening duration of 10 s. (d) Plot of variations in NWs lengths, diameters, and densities with varying V/III BEP ratio. The error bars represent the extrema of the measurements taken from 20 vertical NWs.

decreases the diffusion of Ga and also its desorption from graphene; hence, when simultaneous shutters are opened, the nucleation rate is increased. The droplet shape and the contact angle may thus be more suitable for vertical NW growth. With an increase in temperature, triple phase line of the droplet is maintained within the NW geometry by the impinging group V flux, leading to a continuation of the NW growth. Enhanced Ga desorption at higher growth temperature minimizes parasitic growth. These conjectures are borne out by the observation of complete dewetting of the droplet with the substrate and lack of any parasitic growth for Tg2 of 610 °C (Figure 3d). Using the optimized values of the substrate temperatures Tg1 (540 °C) and Tg2 (580 °C) determined above, we next explored the influence of variation in the V/III BEP ratio (15,

The two temperature process was adopted with the lower temperature used for promoting nucleation while the NW growth occurs predominantly at the higher growth temperature, Tg2. Growth at a single temperature ranging from 520 to 580 °C resulted in nonvertical and crawling NWs with thick parasitic growth (Supporting Information, Figure S1). The success of the two temperature technique for the growth of vertical NWs is consistent with the results reported by Munshi et al.24 for GaAs NWs on graphene. It is to be noted that this is contrary to the NW growth commonly observed on Si substrates where a higher temperature is used for nucleation.2,11,41 The reasons for this difference are unclear; however, we speculate that high Ga diffusion on graphene compared to Si substrate (Figure S2) is the greatest contributing factor for this reversal of growth temperature. Lower temperature D


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Figure 5. SEM micrographs of GaAsSb axial NWs with variation in Ga shutter opening duration of (a) 15 s, (b) 10 s (sample B(20)), (c) 0 s (sample G′(20)), and (d) single NW of 0 s (sample G′(20)). (e) Summary of variations in NW lengths, diameters, and NW densities with varying Ga shutter opening durations. The error bars represent the extrema of the measurements taken from 20 vertical NWs.

Figure 6. (a) TEM image of GaAsSb NW grown on monolayer graphene under optimized conditions. (b−d) HR-TEM images of selected areas marked with white squares in panel (a). (e−g) SAED patterns of selected regions of NW in panel (a), marked with red circles displaying ZB structures viewed from the zone axis of [11̅0]. (h) False color mapping of HAADF image of the corresponding NW. (i−k) EDS line scans taken at the top, middle, and bottom segments of the NW marked with blue squares in panel (h). Scale bar is 75 nm for panel (h).

This optimal ratio was further confirmed by increasing the V/ III BEP ratio to 25, which resulted in a reduced density of NWs (Figure 4d) along with an increase in the aspect ratio from ∼11, correspondings to V/III BEP ratio of 20 to 18 in this case (Figure S3). Finally, the effects of Ga shutter opening duration (15, 10, and 0 s) were investigated by using an optimized V/III BEP ratio of 20. Because the Ga shutter opening is one of the most critical parameters that determine the droplet size during the initial stages of growth, the decrease in NW density (Figure

18 and 20) on NW density with the Ga shutter opening maintained at 10 s prior to growth. Lowering the V/III BEP ratio results in oversized Ga droplets manifesting in shorter NWs with a larger diameter and promoting the formation of horizontal growth, as shown in Figure 4a,b. The NW density was also reduced. A V/III BEP ratio of 20 (Figure 4c) results in a high density of nucleation sites being successfully translated to vertical NW growth due to an abundant supply of group V adatoms consuming the Ga droplet. The ratio of 20 yielded the best results with an optimum Ga droplet shape. E


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ACS Applied Nano Materials 5a) for 15 s duration is attributed to the agglomeration of the droplets, leading to a reduction in the number of droplets being available for nucleation. Lowering the Ga shutter opening from 10 s to simultaneous opening of the shutters resulted in reduced diameter of the NW from 70 to 60 nm and improved vertical NW density from ∼1.6/μm2 to ∼1.8/μm2, as shown in Figures 5b and 5c, respectively. These observations suggest that the droplet size is reduced with reduction in the contact angle, promoting vertical NW density. This, in turn, will lead to variation in the group V species interception with the droplet, hence the degree of supersaturation and Sb incorporation. The graphical representation of the change in length, diameter, and NW density with variation in Ga shutter opening duration is shown in Figure 5d. Continuous parasitic growth was observed along the boundary lines of the graphene, and the region marked with red (Figure S4) suggests NWs were not grown on the non-graphene or oxide areas. AFM tapping mode topography further attests to the growth of vertical nanowires on a larger area of graphene for a Ga opening of 0 s (Figure S5). Figure 6a shows a TEM image of a NW that was grown under optimized conditions (sample G′(20)). HRTEM images (Figure 6b−d) and selected area electron diffraction (SAED) patterns (Figure 6e−g) confirm the zinc-blende (ZB) structure of the GaAsSb NWs. Microtwins and stacking faults were observed only near the NW tip as shown in Figure 6b with the presence of microtwins being identified by the faint satellite spots observed in the corresponding SAED pattern (Figure 6e). The occurrence of these structural irregularities at the tip, commonly observed in GaAsSb NWs of low Sb composition,15 is attributed to the concurrent closing of all the shutters during the termination of the Ga droplet, which yields an As-rich Ga droplet with the wurtzite structure of the NW being more favored, leading to creation of these planar defects. False-color high-angle annular dark-field (HAADF) TEM imaging (Figure 6h) and the corresponding EDS compositional line scans (Figure 6i−k) reveal uniform compositional homogeneity along the NW with an Sb composition of 7.3 at. %. The intensity of the counts for this EDS mapping is not sufficiently high to enable quantitative mapping. Hence, any composition noted on the line scans would deviate significantly from the real value. The 7.3 at. % Sb is a more accurate representation of the real Sb composition since it was obtained from the integrated signals from the entire mapping. In addition, the at. % at the two sides of the line scan (that is, in the vacuum area) is an artifact because the software attempts a compositional analysis on the vacuum area adjacent to the nanowire. The HRTEM image and EDS compositional line scans of another shorter NW from sample G′(20) (Figure S6) reveal triple phase line of the droplet pinned at a top facet of the NW, which has droplet geometry not completely covering the tip of the NW. The droplet offset to the side is the origin of horizontal or crawling NWs grown on graphene.30 Hence, we believe the TEM shown in Figure 6a, which has round top, indicates that the Ga droplet completely covering the NW tip leads to vertical nanowires. X-ray diffraction spectra on NWs for three different V/III BEP ratios under simultaneous opening of Ga, As, and Sb shutters (sample G′(20)) and with a 10 s Ga opening, V/III = 20 (sample B(20)) reveal the existence of only GaAsSb (111) and Si (111) Bragg peaks as shown in Figure 7a, attesting to the vertical alignment of NWs. The shift in the GaAsSb (111) Bragg peak with respect to that of GaAs1−xSbx NWs (similar Sb

Figure 7. (a) XRD spectra of GaAs1−xSbx NWs on graphene for the V/III BEP ratios of 20, 18, and 15, corresponding to the growth conditions of the simultaneous opening of all shutters and V/III BEP ratio of 20 for Ga shutter opening duration of 10 s. For comparison, a reference XRD spectrum of GaAsSb NWs of similar Sb composition grown on Si is also shown. (b) The corresponding GaAsSb (111) Bragg peak positions of the XRD spectrum are shown.

composition) grown on Si with the same V/III BEP ratio shifts toward a lower angle with a decreased Ga opening duration as well as with a decreasing V/III BEP ratio (Figure 7b), indicating increased Sb alloying. Thus, the optimal growth conditions of V/III BEP ratio of 20 with simultaneous opening of shutters were found to lead to the highest Sb incorporation compared to other growth conditions. The contribution of parasitic growth to the XRD spectra is ruled out by comparing exclusively with the XRD of the parasitic growth with no NWs as shown in Figure S7. The 4 K PL spectra of SNW grown under optimized growth conditions corresponding to a V/III BEP ratio of 20 with simultaneous opening of all shutters (sample G′(20)) exhibited a well-defined sharp PL peak energy at 1.32 eV with a fwhm of 46 meV, while a NW grown at a V/III BEP ratio of 20 and a Ga shutter opening duration of 10 s (sample B(20)) exhibited a PL energy peak at 1.35 eV with an fwhm of 69 meV, as illustrated in Figure 8a. The sharp peak of the PL spectra with simultaneous opening of all the shutters indicates compositional homogeneity of NWs grown via smaller Ga droplets. The PL spectra exhibited a red-shift of ∼137 meV with a decrease in the V/III BEP ratio from 20 to 15 irrespective of the Ga shutter opening. Such large red-shifts with a change in V/III BEP ratios are not observed when the NWs are grown on a Si substrate.2 The Sb composition based on the literature data for thin films42 indicates a variation from 5.4 to 6.5 at. % for varying the V/III BEP ratio from 15 to 20 for simultaneous opening of the shutters, while it shows a larger variation from 1.2 to 7.5 at. % for a similar V/III BEP ratio of Ga shutter opening duration of 10 s. These large redshifts observed on graphene are associated with Sb-induced modulation of the droplet size, as explained below. First, the migration energy of Sb adatoms (0.03 eV)26 is much smaller than that of As adatoms (0.2 eV) on graphene. Furthermore, Sb, which acts as a surfactant, impacts the migration length of other adatoms, namely Ga,11,16,43 and also lowers the contact angle of the Ga droplet30 due to the lowering of the surface energy of Ga droplet.16,30 The effect of lower migration energy of Sb and its nature of riding on the surface due to its surfactant effect likely enhances the probability of interaction with Ga, favoring Ga−Sb bonding more than Ga−As bonding, which would promote higher incorporation of Sb in the NWs. The reduction in the contact angle due to Sb also has a favorable influence on the droplet shape, which otherwise would result in a large contact angle on the graphene44 due to the low surface energy of the substrate,30 making it unsuitable F


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Figure 8. (a) Low-temperature (4 K) normalized PL spectra of GaAsSb SNW for Ga shutter opening duration of simultaneous opening of all shutters (samples G′(20), H′(18), and I′(15)) and 10 s (samples B(20), E(18), and F(15)) for corresponding V/III BEP ratios of 15, 18, and 20. (b) Comparison of 4 K μ-PL spectra of single NW grown on graphene and silicon and (c) plot of full width at half-maximum versus V/III BEP ratio for Ga shutter opening durations of 0 and 10 s.

III−V adatoms adsorption18,26 on graphene, as opposed to (111) Si with no preferential sites. Furthermore, the lattice mismatch of GaAs0.9Sb0.1 with graphene is ∼6.01% and ∼8.53% for the H site and both H, B sites, respectively. This is larger than the ∼4.4% mismatch on (111) Si substrate having an adverse impact on the NW density on graphene. Another factor that likely contributes to the low NW density is the significantly lower surface energy of graphene compared to that of Si, as mentioned earlier. A comparison of the 4 K PL spectra of the GaAsSb SNW grown on graphene with those grown on Si (Figure 8b) for similar PL peak energy grown under optimized conditions (optimized conditions being different from NWs grown on Si) shows that the NWs grown on graphene (sample G′(20)) exhibit a 3-fold higher intensity with a lower fwhm of 46 meV as opposed to 152 meV for those nanowires grown on Si (Figure 8c). The higher PL intensity observed in NWs grown on graphene is speculated to be due to better compositional homogeneity as well as a reduction in point defects as explained in the following discussion. The low migration energy, relatively smooth surface presented by graphene due to lack of significant parasitic layer, and the diminished surfactant effect of Sb on reducing the adatom mobility of Ga likely contribute to the relatively high Ga diffusion on graphene. This leads to a reduction in the concentration of point defects, namely, the Ga antisite/Ga vacancy complex defect commonly present in these alloys. Furthermore, the low NW density also reduces the shadowing effect from neighboring NWs and As− Sb exchange possibility due to re-emission from the neighboring side facets leading to less compositional variation in the NWs. The EDS data show evidence of uniform composition as opposed to our earlier work on Si where, particularly at the bottom segment, there is a pronounced effect of As−Sb exchange.2,48

for vertical NW growth. Glancing angle SEM images of Ga droplets on Si and graphene substrate reveal contact angles to be ∼90° and ∼120°, respectively. The droplet on the graphene exhibits an asymmetric shape (Figure S8). Also, GaAsSb droplets before the nucleation of NW for a 60 s shutter opening duration at a growth temperature of 540 °C suggest that the desired droplet contact angle for vertical NW growth was ∼120°. Furthermore, opening all the shutters simultaneously is the optimized condition, as it assists not only in the Sb-induced reshaping of the droplet to establish a favorable contact angle for growth but also in rapid consumption of the droplet, promoting successful growths of smaller droplets with high growth rates. High Ga adatom mobility on the graphene likewise contributes to the observed high growth rate, smaller droplets, and diminished nucleation sites44 as compared to those commonly observed on Si for similar Sb composition.29 The above discussion is also consistent with other subtle differences observed between the NWs grown on graphene and Si. These include the Ga opening for a few seconds being essential for the growth of Ga-assisted self-catalyzed NWs,1,2,5,11,37,45−47 thicker NWs of similar Sb composition, and lower growth rates for the GaAsSb NWs on Si. Furthermore, the above surface engineering via Sb also explains the higher vertical NW density that is obtained with simultaneous opening of the shutter conditions as well as somewhat higher (∼2×) vertical NW density (1.8/μm2) in the GaAsSb NWs as compared to the GaAs NWs (1/μm2) grown on graphene reported by Munshi et al.24 This trend in NW density with Sb incorporation is also in contrast to with those commonly observed on GaAsSb NWs grown on a Si substrate where Sb incorporation adversely affects the NW density, which further attests to the effect of the Sb surfactant effect. It is to be noted that NW density is significantly lower than those grown on (111) Si (Figure S9). This is attributed to only two of the three preferred H and B sites being favorable for G


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ACS Applied Nano Materials Room temperature (RT) Raman spectra (Figure 9) of optimized GaAsSb NW growth on graphene exhibited two

Figure 9. (a) RT Raman spectra of GaAsSb NWs for an optimized sample and Ga shutter opening duration of 10 s, grown on graphene. Raman spectra of (b) graphene before and after GaAsSb NW growth (sample G′(20)).

Figure 10. I−V characteristics of GaAsSb single NW grown on graphene layer and Si substrates in the dark, under illumination with a 673 nm AFM laser and 633 nm visible laser. The inset shows the dark current for single NW grown on Si and graphene substrates. The finite element modeled nanowire of the corresponding sample simulated by COMSOL Multiphysics software data is indicated as dots which are superimposed on the experimental plots.

peaks at 260.7 and 282.2 cm−1, which correspond to the transverse optical and longitudinal optical phonon modes, respectively.2 An additional peak was also observed for both optimized growth conditions (sample G′(20)) and Ga opening duration of 10 s (sample B(20)) around 230 cm−1, which correspond to a GaSb-like TO mode,2,49 which is marked with an asterisk in Figure 9a. This is commonly observed at higher Sb concentrations in the GaAs matrix.2 The observation of graphene-related G and 2D Raman peaks at 1581 and 2674 cm−1, respectively, and the ratio of their peaks I2D/IG > 1 in Raman spectra on graphene samples before and after growth indicate monolayer graphene underneath the NWs, as shown in Figure 9b. The presence of a G band suggests covalently bonded carbon atoms participating in the in-plane C−C stretching vibration. This has also been viewed as strong evidence for the presence of graphene in the literature, for instance, growth of ZnO nanorods on graphene by Das et al.50 Furthermore, the Raman peak intensity at 1350 cm−1 related to the D mode of graphene near the noise level in both of these samples attests to minimal defects in the graphene layer. Finally, as noted in the Experimental Details section, the CSAFM technique was used to measure the I−V characteristics51 of a single vertical GaAsSb NW grown on graphene as well as on Si. The NWs grown on graphene exhibited photoresponse of few nA/W to few μA/W depending on the bias voltage and illumination power as shown in Figure 10. These were consistently close to an order of magnitude higher than those observed for NWs grown on Si and also exhibited significantly lower dark current in the picoamperes range (Figure 10, inset). The best fit of a simulated I−V curve to experimental data under dark conditions for NWs grown on graphene results in an electron (μn) mobility of 182 cm2 V−1 s−1, hole mobility (μp) of 53 cm2 V−1 s−1, and a carrier concentration of 5.3 × 1014 cm−3. The results on Si are marginally lower with μn of 173 cm2 V−1 s−1 and μp of 47 cm2 V−1 s−1 at a background carrier concentration of 2.5 × 1015 cm−3. Under 30 mW laser illumination, the carrier concentration increased in both cases to ∼7.3 × 1015 and 3.1 × 1015 cm−3 for the NWs grown on graphene and Si, respectively. The carrier mobilities under illumination also improved in both cases, but the NWs on graphene still exhibited higher μn and μp values (248 and 71 cm2/(V s), respectively), over the NWs on Si (with corresponding values of 195 and 55 cm2/(V s)). The improved characteristics under illumination in NWs grown on graphene are therefore due predominantly to the lower

background carrier concentration and a slightly enhanced mobility. This investigation clearly demonstrates the growth of highquality GaAsSb NWs on graphene via the use of Sb as a surfactant for surface engineering of the graphene. The pronounced effect of the V/III BEP ratio on shifting the PL peak is very attractive because it shows a pathway for realizing photodetectors with improved performance in the telecommunication wavelength regime between 1.3 and 1.55 μm based on single NWs with a higher Sb composition. Potential pathways to further improve the vertical GaAsSb NW density on graphene are by enhancing the wetting of the graphene surface using different surface treatments or/and use of an appropriate surfactant during the initiation of growth that would modulate the contact angle of the droplet as well as limit the Ga adatom diffusion.

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CONCLUSION The two temperature growth steps along with other growth optimization parameters and the surfactant effect of Sb were successfully used to realize GaAsSb NWs with a vertical orientation and reasonable density on monolayer graphene. Lowering the V/III BEP ratio leads to higher Sb incorporation in the NWs. GaAs1−xSbx NWs grown under optimal conditions resulted in a sharp 4 K PL emission peak at 0.94 μm (1.319 eV) with an fwhm of 46 meV and ∼3-fold higher intensity as well as higher optical responsivity predominantly due to a lower background carrier concentration with somewhat improved carrier mobilities compared to the NWs grown on a Si substrate, illustrating the high optical quality of the NWs. This conclusion is also borne out by the I−V data, which revealed dark current significantly lower for NWs grown on graphene in comparison to those grown on Si. This shows the potential for a NW-based photodetector grown on monolayer graphene for achieving the superior figure of merit and fabrication of a flexible NIR photodetector.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsanm.9b00893. H


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Variation of single growth temperature of GaAsSb NW growth on graphene, Ga diffusion on Si and graphene/ SiO2/p-Si (111) substrates, GaAsSb NW growth for sample J′(25) growth conditions, parasitic growth on boundaries of graphene, comparison of SEM and AFM images of GaAsSb NWs on graphene at similar magnification, TEM images of shorter NW of sample G′(20), XRD spectra of GaAsSb parasitic growth on graphene/SiO2/p-Si (111) substrate, contact angles of Ga droplets on Si and graphene/SiO2/p-Si (111) substrates, morphology of GaAsSb NWs grown on p-Si (111) substrate (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Surya Nalamati: 0000-0002-2180-1769 Shanthi Iyer: 0000-0002-8163-9943 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This material is based upon research supported by the Office of Naval Research under Award N00014-16-1-2720. Part of this work was performed at the Joint School of Nanoscience and Nanoengineering, a member of the Southeastern Nanotechnology Infrastructure Corridor (SENIC) and National Nanotechnology Coordinated Infrastructure (NNCI), which is supported by the National Science Foundation (ECCS1542174). The authors acknowledge the use of the Analytical Instrumentation Facility (AIF) at North Carolina State University, which is supported by the State of North Carolina and the National Science Foundation (Award ECCS1542015). The AIF is a member of the North Carolina Research Triangle Nanotechnology Network (RTNN), a site in the National Nanotechnology Coordinated Infrastructure (NNCI).

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