Growth of GaN Micro- and Nanorods on Graphene-Covered Sapphire

Article pubs.acs.org/crystal

Growth of GaN Micro- and Nanorods on Graphene-Covered Sapphire: Enabling Conductivity to Semiconductor Nanostructures on Insulating Substrates Martin Heilmann,*,† George Sarau,† Manuela Göbelt,† Michael Latzel,†,‡ Sumesh Sadhujan,† Christian Tessarek,†,‡ and Silke Christiansen†,∥ †

Max Planck Institute for the Science of Light, Günther-Scharowsky-Str. 1, 91058 Erlangen, Germany Friedrich-Alexander-Universität Erlangen-Nürnberg (FAU), Institute of Optics, Information and Photonics, Staudstr. 7/B2, 91058 Erlangen, Germany ∥ Helmholtz Centre Berlin for Materials and Energy, Hahn-Meitner Platz 1, 14109 Berlin, Germany ‡

ABSTRACT: The self-catalyzed growth of vertically aligned and hexagonally shaped GaN micro- and nanorods on graphene transferred onto sapphire is achieved through metal−organic vapor phase epitaxy. However, a great influence of the underlying substrate is evident, since vertically aligned structures with a regular shape could not be grown on graphene transferred to SiO2. The optical properties of the regular GaN nanorods were investigated by spatially and spectrally resolved cathodoluminescence showing defect related emission only near the interface between the sapphire substrate and nanorods but not from their upper part. Micro-raman spectroscopy confirms that the single-layer graphene remains virtually unchanged in terms of the Raman signal, even after undergoing high temperatures (∼1200 °C) during nanorod growth. Furthermore, Raman mapping demonstrates that GaN structures predominantly grow on defective parts of graphene, giving new insight into the nucleation and growth mechanism of semiconductors on graphene. To validate the conductivity of graphene, when being attached to the sapphire substrate and after the nanorod growth, current−voltage investigations were carried out on single, as-grown, GaN nanorods with a nanoprober in a scanning electron microscope. These measurements demonstrate the viability of graphene as a conductive electrode, for example, as a back contact for GaN nanorods grown on insulating sapphire.

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INTRODUCTION GaN and related alloys (AlGaN, InGaN), widely used in optoelectronic or high-power electronic devices, are usually grown heteroepitaxially with a certain lattice mismatch to substrates such as sapphire, SiC, or more recently Si.1−3 To avoid substantial plastic relaxation of the lattice mismatch induced strain by the formation of extended defects such as dislocations or stacking faults, it is beneficial to grow nanorods (NRs) instead of layers.4−7 Thereby, one can exploit a two- or three-dimensionally limited growth, which promotes elastic rather than plastic strain relaxation at the free surfaces that extends in the volume of the NRs. In NRs, already used in device concepts such as light-emitting diodes or solar cells, elastic strain relaxation together with a reduction of performance-limiting defects at the free surfaces and in the bulk have been proven to result in virtually stress-free material.7−11 Whispering gallery modes (WGMs) as well as lasing activity in GaN micro- and nanorods grown self-catalyzed on sapphire have been demonstrated by cathodo- and photoluminescence.7,12−14 The existence of WGMs are indicative of a good morphological, structural, and ultimately optical quality of the rod material, making these structures promising for integration in light emitting devices (e.g., diodes, lasers, and optical © XXXX American Chemical Society

sensors). However, the insulating nature of the sapphire substrate used for the epitaxial growth is disadvantageous for the implementation of those micro- and nanorods in optoelectronic device concepts in which bottom and top electrical contacts are mostly required. Recently, graphene with its two-dimensional planar configuration of sp2-bonded carbon atoms has attracted great interest not only as a transparent top contact material but also as a growth substrate for semiconductor nanowires.15 The stability of graphene at high temperatures, its flexibility and optical transparency as well as its excellent electrical conductivity makes it a viable substrate for the growth of GaN NRs.16 Even if graphene is transferred to the transparent but insulating sapphire, it can serve as a back contact in this configuration. Today’s capacity to synthesize large area graphene layers by chemical vapor deposition (CVD) is also useful for its implementation in optoelectronic devices.17 It was already shown by Munshi et al.18 that in theory, epitaxial growth of GaN on pristine graphene should be possible with adsorption sites within the hexagonal carbon rings and between Received: October 13, 2014 Revised: January 30, 2015

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DOI: 10.1021/cg5015219 Cryst. Growth Des. XXXX, XXX, XXX−XXX

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Crystal Growth & Design

For the growth of n-doped GaN NRs, the graphene/sapphire substrates were loaded in an Aixtron 200RF horizontal flow MOVPE reactor. The growth temperature was monitored by a thermocouple inside the susceptor. It is estimated that the surface temperature of the substrate on the susceptor is approximately 50 to 100 °C lower than the value measured by the thermocouple. To improve the homogeneity of the growth process across the sample, the substrate was rotated through a gas foil rotating holder using nitrogen (N2) with a flux of 100 sccm. All process steps were carried out at a pressure of 100 mbar. After inserting the sample, the reactor was evacuated and afterward purged with N2 in order to remove any trace of oxygen or water which cannot only be detrimental to the graphene at elevated temperatures but also can react violently with the precursors used during the growth. In a first step, the substrate was then thermally cleaned of contaminants such as PMMA residues in hydrogen (H2) atmosphere at 1200 °C for 5 min. Ammonia (NH3) was subsequently introduced into the reaction chamber for the nitridation of the substrate surface with a flow of 1500 sccm for 10 min under H2 atmosphere. On plain sapphire substrates without graphene this step forms a thin N-polar AlN layer on the surface.27 The induced polarity of the substrate is a crucial parameter, which determines whether rodlike or pyramidal island growth occurs.28 Afterward, a GaN nucleation layer was deposited on the substrate for 16 s by introducing trimethylgallium (TMGa) with a flux of 180 μmol/min and NH3 with a flux of 1500 sccm at 1200 °C under H2 atmosphere. GaN micro- and nanorods were grown on this nucleation layer after a second nitridation. To this end, TMGa with a flux of 45 μmol/min and NH3 with a flux of 25 sccm were introduced into the reactor chamber together with diluted silane (SiH4) as an n-type dopant at 1150 °C for 13 min. Here SiH4 has a strong influence on the vertical growth of the NRs as pointed out and quantified in the literature.29 After finishing the growth of the rods, the NH3 flux was kept constant to stabilize the GaN structures until the reactor was cooled down to room temperature.7 SEM images were acquired with a Hitachi S4800 SEM at an acceleration voltage of 5 keV at an angle of 60° between the electron beam and surface normal. The SEM is also equipped with a CL measurement unit (Gatan MonoCL), which was used for the optical characterization of single nanorods. The CL measurements were performed at room temperature under the same SEM settings as for imaging the structures. Micro-Raman measurements were carried out at roomtemperature in the backscattering configuration using a LabRam HR800 spectrometer from Horiba Scientific. Two linearly polarized lasers emitting at 457 and 660 nm were employed for the Raman excitation. The laser light was tightly focused by a 100× objective (numerical aperture 0.9), resulting in a diameter of the normally incident probing beam of ∼0.7 μm and ∼1 μm corresponding to a laser power on the sample surface of ∼582 and ∼1.28 mW using filters. These laser powers ensured no structural changes to the graphene and the GaN NRs as well as no Raman shift due to local heating by the laser beam. To study the electrical conductivity between GaN NRs and graphene, an integrated-circuit-testing device with motorized probing arms (Kammrath & Weiss GmbH) inside a SEM (FEI) was used to contact single NRs with tungsten nanoprobes and measure the current−voltage (I−V) characteristics via a semiconductor characterization system (Keithley SCS 4200).

two carbon atoms. In spite of these advantages of graphene, it has been stated that only weak nucleation and thus slow, island layer growth rather than two- or three-dimensional epitaxy is expected if semiconductor structures are grown on graphene. This is supposed to be the result of high surface tension caused by the lack of dangling bonds in pristine, nonfunctionalized graphene layers.18 Such a growth behavior is unwanted since controlled layer or NR growth strongly relies on epitaxy that establishes the growth direction of NRs reproducibly. So far, GaN microdiscs or nanostructures were realized by epitaxial lateral overgrowth of intermediate ZnO nanowalls or nanowires on graphene.19,20 Also AlN, which can act as a nucleation or buffer layer, was grown epitaxially on graphene in a first step toward producing GaN structures.21 Furthermore, it has been shown that GaN layers or irregular micro- and nanostructures can be obtained epitaxially on single- or multilayered graphene with either sapphire or SiO2 as the initial substrate.22−25 In the latter case, vertical aligned microstructures were grown through an intermediate GaN buffer layer on the graphene. GaN layers were also grown on epitaxial graphene with SiC as the underlying substrate, and a functional and even flexible LED structure was shown after the release of the layers from the initial substrate.26 In the present paper, we show that vertically aligned and hexagonally shaped n-doped GaN micro- and nanorods can be grown self-catalyzed via metal−organic vapor phase epitaxy (MOVPE) directly on single-layer graphene transferred onto sapphire substrates without any intermediate layer. Cathodoluminescence (CL) measurements indicated the good optical and structural properties of the GaN NRs. Micro-Raman spectroscopy proved that graphene is resilient against the growth conditions during the MOVPE and confirmed the presence of graphene underneath the GaN NRs. Furthermore, correlated micro-Raman mapping and scanning electron microscopy (SEM) imaging demonstrated that GaN NRs grew predominantly on defective regions of graphene, giving new insights into the nucleation and growth mechanisms of GaN micro- and nanostructures on graphene. The electrical conductivity between single, as-grown GaN NRs and graphene as the bottom contact was studied using a nanoprober setup within a SEM. Finally, single GaN NRs were bent with a nanoprober until their fracture in order to investigate the bonding between the NRs and the sapphire as well as the graphene/sapphire substrates. It is implied that the NRs nucleate and grow both on graphene and on sapphire through nanoholes in graphene.

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EXPERIMENTAL SECTION The substrates for the growth of GaN NRs were prepared by transferring predominantly single-layer CVD graphene onto csapphire. The graphene films were grown on Cu foils by a twostep CVD process at a temperature of 1050 °C and a pressure of 1.33 mbar following procedures from the literature.17 After the growth, polymethylmethacrylat (PMMA) was spin-coated on graphene as a protective and stabilizing layer during the graphene transfer onto the substrate of choice. The Cu foil was subsequently etched away in an aqueous solution of ammonium persulfate (2.5% wt) to release the graphene/PMMA film. This film sandwich of ∼1 cm2 was transferred onto a sapphire substrate, and the PMMA was dissolved in a 1:1 mixture of acetone and 1,2-dichloroethane, leaving only the graphene layer on the sapphire substrate. B


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nucleation layer, more supplied material was accumulated at these positions in the subsequent growth of NRs leaving large areas of pristine graphene uncovered. Therefore, micro- and nanorods were obtained, depending on the local defect configuration of graphene, while the rods were larger as compared to the growth on the reference sample where there were more nucleation points present. A correlation between the density of nucleation points and the height and diameters of GaN structures grown on them was already reported for the self-catalyzed growth on sapphire.7 To investigate the influence of the underlying substrate on the growth of GaN NRs, graphene was also transferred onto amorphous SiO2, and the same growth procedure as described above was applied. SiO2 is commonly used as a masking material (e.g., for selective area growth of GaN NRs) and essentially does not permit any growth of GaN.30 The SEM image in Figure 2a shows nearly no growth of GaN on bare

Moreover, the tungsten nanoprobes were employed to bend and remove single nanorods from the substrates.

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RESULTS AND DISCUSSION Figure 1a shows a SEM image of GaN micro- and nanorods on a graphene/sapphire substrate, which were grown by the aforementioned process. Hexagonally shaped and vertically aligned GaN structures with varying aspect ratios are visible. Figure 1b displays a close-up view of a hexagonal GaN NR with well-defined edges as well as straight and smooth sidewall facets. The graphene layer has no influence on the vertical alignment and the hexagonal shape of the structures, if compared to a reference sample (Figure 1c) grown on sapphire without graphene under the same growth conditions. However, we obtained a higher uniformity and density of NRs on the reference sample (∼1 × 108 cm−2, Figure 1c) as compared to the sample with graphene (∼6 × 106 cm−2, Figure 1a). Moreover, the structures on graphene (mean diameter of ∼500 nm and an average height of ∼4 μm) are considerably larger than those on the reference sample (mean diameter of ∼250 nm and an average height of ∼1.5 μm). A close-up view of GaN NRs on sapphire can be seen in Figure 1d, where the hexagonal shape of the structures is clearly visible. It is expected that less GaN nucleation sites were formed on the surface of graphene during the initial deposition of the nucleation layer due to the lack of chemical reactivity of graphene leading to the lower density of NRs as compared to the epitaxial growth directly on sapphire. The spatial distribution and density of GaN NRs is related to the presence of defects in graphene. These defects provided the dangling bonds necessary to initiate the growth of a nucleation layer at the graphene surface that consequently enabled the epitaxial growth of the GaN NRs. This growth mechanism is in line with that suggested for the growth of GaAs nanowires on graphitic surfaces.18 Since GaN mainly nucleates at defects on the graphene during the growth of the

Figure 2. (a) SEM image of GaN structures grown on graphene/SiO2 substrates. The bare SiO2 remains mostly free from GaN, while on the graphene, a high density of GaN microstructures can be observed. (b) A close-up view on the hexagonal GaN microstructures shows no preferential but rather a random growth direction. Scale bars in (a) and (b) are 50 and 10 μm, respectively.

SiO2, while the graphene is densely covered with GaN microstructures. It is expected that the nucleation of GaN mainly occurred at defects in graphene since pristine graphene and SiO2 hinder the epitaxial growth. A close-up view in Figure 2b shows the irregularly shaped GaN microstructures, which are comparable to those reported by Chung et al.25 grown on graphene transferred to SiO2 substrates as well. However, no aligned growth occurs as compared to the growth on graphene/ sapphire substrates. These results emphasize the strong influence of the initial substrate underneath the single-layer graphene on the growth of semiconductor nanostructures on top of transferred graphene.

Figure 1. (a) SEM images of GaN micro- and nanostructures grown on the graphene/sapphire substrate and (b) a close-up view of a GaN NR. (c) As a reference, the same growth parameters were applied to a sapphire substrate without graphene, resulting in a denser NR structure and (d) a close-up view of GaN NRs on sapphire. The scale bars are 10 μm in (a and c) and 2 μm in (b and d). C


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Crystal Growth & Design In order to investigate the optical properties of the GaN NRs at their top and near the interface to the substrate, CL measurements were conducted at single, as-grown NRs. In Figure 3a, at the left side a SEM image of a GaN NR grown on graphene can be seen with the black and blue crosses, indicating where the electron beam was focused during the acquisition of the CL spectra in Figure 3c. The panchromatic CL image of the same NR on the right side of Figure 3a displays a homogeneous intensity in the upper part, while near the interface to the graphene/sapphire substrate a higher intensity is evident, where one facet at the base also appears brighter than the other. This is different to GaN NRs grown on a sapphire substrate, as seen in the SEM and CL images in Figure 3b. The red and green crosses again depict were the electron beam was focused during the acquisition of the CL spectra. In the panchromatic CL map of this sample area in Figure 3b, on the right side, the base of the GaN NRs appears less bright, while a homogeneous emission can be seen in the upper part of the NR. In the spectra in Figure 3c, only the near band edge emission of GaN at 365 nm was detected at the top of the NRs (black and red curve), while in the spectrum acquired at the base of the NR, an additional broad defect related band centered around 560 nm became prominent. It has been proven that this so-called yellow defect luminescence is originating from interband transitions, which can occur at high defect densities generated to accommodate the lattice mismatch between the GaN NRs and sapphire close to their interface.31 However, the intensity of the yellow defect luminescence appears higher at the base of the GaN NRs grown on graphene, which can be attributed to carbon impurities as shown in the literature.32 Carbon as a contamination source might be originating from PMMA residues on defects in graphene, which were not completely removed during the initial chemical and thermal cleaning of the substrate, or from carbon atoms that were extracted from edges or defects in the graphene. The brighter appearance of one facet at the base of the NR in Figure 3a at the right side was observed in numerous other NRs on graphene. This can be attributed to a combination of locally different carbon incorporation and point defects as well as an inversion domain inside the rod.32 Further investigation has to be conducted in order to separate these effects and clarify the observed differences. Nonetheless, these CL measurements are indicative of defect-free GaN material in the upper part of the NRs resulting in good optical as well as structural properties comparable to those of similar structures grown directly on sapphire.7 Figure 4 summarizes the Raman results obtained with the 457 nm (Figure 4, panels a and d) and the 660 nm (Figure 4, panels b and e) lasers focused on graphene at a representative position shown in the SEM image (Figure 4c). In addition to the GaN microrod, one can see light and dark SEM contrasts related to virtually defect-free and defective areas of graphene, respectively, as proven in the following by the Raman data. The first important point to be addressed is the influence of the MOVPE (i.e., high temperatures with various precursor gases on the graphene itself). Typical Raman spectra of graphene before (black line) and after (red line) the growth of GaN NRs, measured apart from the GaN rod in an area of light SEM contrast are displayed in Figure 4a. The presence of the G and 2D peaks and the absence of a clear D peak in both spectra demonstrates that graphene withstood the growth conditions for GaN NRs accompanied by negligible defect formation, since for the Raman activation of the D peak at ∼1350 cm−1 a defect

Figure 3. (a) SEM image of a GaN NR grown on the graphene/ sapphire substrate (left), panchromatic CL map of the same region (right). (b) SEM image of GaN NRs grown directly on the sapphire substrate (left), panchromatic CL map of the same region (right). (c) Room temperature CL spectra taken with a fixed electron beam at the base and at the top of GaN NRs obtained with and without graphene on sapphire. The colored crosses in (a and b) indicate where the electron beam was focused during the acquisition of the CL spectra with the respective colors in (c). The spectra from the base and the top are vertically shifted for clarity. The scale bars in (a and b) are 2 μm.

in graphene is required.33,34 Furthermore, both spectra show a symmetric line-shape of the 2D-peak, a 2D full width at halfmaximum of ∼25 cm−1, and an intensity ratio between the 2D and G peaks > 2, which are characteristic for monolayer D


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Figure 4. (a) Raman spectra of single layer graphene on sapphire taken with an excitation wavelength of 457 nm before (black line) and after (red line) the MOVPE on an area of light SEM contrast as well as in the vicinity of the GaN microrod on an area of dark SEM contrast (blue line) in (c). (b) By using a wavelength of 660 nm, the graphene Raman peaks were detected through the GaN microstructure (blue line) as verified by similar measurements on a microrod grown directly on sapphire (red line) and on the graphene without microrod (black line). The spectra in (a) and (b) are vertically shifted for clarity. (c) SEM image of the sample area where the Raman spectra and maps were acquired with a large GaN microstructure chosen for the Raman mapping. (d) Map of the D-peak intensity measured with an excitation wavelength of 457 nm, depicting the defective regions in graphene. (e) Map of the 2D-peak intensity excited by a wavelength of 660 nm showing that graphene withstood the MOVPE growth conditions even below the GaN microstructure. The white circle represents the position of the GaN microstructure from (c).The scale bars in (c, d, and e) are 5 μm.

graphene, while the 2D-peak shift is attributed to nitrogen doping of graphene during the MOVPE process.35−37 For comparison, we display a typical Raman spectrum acquired after the growth next to the GaN microrod in a region of dark SEM contrast (blue line in Figure 4a). Distinct D and D′ peaks are visible along with the background, originating from the yellow defect luminescence of GaN (Figure 3c, blue line). It should be noted that although the incident laser light was focused on the graphene, the GaN microrod was excited by the light cone when measuring in its vicinity giving rise to the observed background. The second critical aspect is the confirmation of graphene beneath the GaN structures, that is, the Raman probing of graphene through the GaN micro- and nanorods. Because the 457 nm laser along with the excited Raman light from the GaN rod itself [A1(TO), E1(TO), E2(high) at ∼468.4, 469, 469.2 nm] and the underlying graphene (D, G, 2D at ∼487.6, 493.1, 522.2 nm) strongly pumped the yellow defect luminescence in GaN (left-side on the increasing part of the yellow defect luminescence in Figure 3c, blue line), the Raman peaks of graphene were hardly noticeable on top of the rods being covered by this emission, while increasing the acquisition time led quickly to the detector saturation. The situation was improved for the 660 nm laser as evidenced by the representative Raman spectra shown in Figure 4b, where the Raman peaks of graphene (black line, measured in a defective area away from the rod) were clearly detectable when probing on top of/beneath the microrod (blue line) despite the background emerging from the defects in the GaN microstructure as confirmed by similar measurements on a microrod

grown directly on sapphire (red line). Compared with the 457 nm laser, the background was less pronounced because both the 660 nm laser and the stimulated Raman light from the GaN rod itself [A1(TO), E1(TO), E2(high) at ∼683.1, 684.4, 684.9 nm] and the graphene underneath the GaN rod (D, G, 2D at ∼722.5, 737.1, 798.5 nm) pumped the yellow defect luminescence of GaN only weakly (right side on the decreasing part of the defect emission in Figure 3c, blue line). The experimental approach established herein can be extended to other nano-object/graphene heterostructures (e.g., optical studies of their interfaces that require the minimization of intrinsic and/or defect-induced light emission from the nanoobject). Figure 4 (panels d and e) show spatially resolved Raman maps (step size 0.5 μm) measured on the area displayed in the SEM image in Figure 4c. The large GaN microrod (∼4 μm diameter), suitable for Raman mapping, grew on a line exhibiting a dark SEM contrast similar to the GaN NR in Figure 3a. An excitation wavelength of 457 nm was used to generate the D-peak intensity map in Figure 4d. The good correlation between the regions with nonzero intensity of the D-peak and dark SEM contrast proves the presence of a high defect density at these positions. The defective graphene can provide dangling bonds for the nucleation, diffusion, and growth of GaN micro- and nanorods. Thus, the defect-free areas of graphene remain mostly free from GaN structures, which can explain the lower density of NRs on graphene/ sapphire as compared with the bare sapphire substrates, which offers more nucleation points (see Figure 1 for comparison). Such defective lines may evolve during the CVD growth of E


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Crystal Growth & Design graphene when grains, nucleated at different locations, grew laterally and finally coalesced, leading to locally incomplete hexagonal carbon structures.38,39 Due to the large spot size of the excitation lasers, such nanoholes in the graphene cannot be detected within the defective lines, but they contribute to the overall D-peak intensity. It is worth noting here that the intensity of the D-peak line below the GaN microstructure, highlighted by the white circle in Figure 4 (panels d and e), is not zero but the D-peak is simply covered by the high yellow defect luminescence of GaN. A 660 nm laser was employed to generate the 2D-peak intensity map in Figure 4e. Since the 2Dpeak does not need a defect for its activation, it can be detected over the whole sample area and from underneath the GaN microstructure due to the proper choice of the excitation wavelength. The variations in the 2D-peak intensity can be attributed to intrinsic differences in the CVD graphene layer. This finding demonstrates that graphene withstands the MOVPE under the GaN structures, making it a powerful candidate in terms of mechanical, thermal, and structural integrity for the growth of semiconductor nanostructures. For demonstrating the electrical conductivity between ndoped GaN NRs and graphene beneath them, tungsten nanoprobes were used to contact the as-grown NRs as schematically illustrated in Figure 5a. At first, two tungsten probes were placed on graphene, at a distance of 30 μm apart from each other. The resulting symmetric current−voltage (I− V) characteristic can be seen in the semilogarithmic plot in Figure 5c (blue curve), indicating a good contact between tungsten and graphene. In the next step, an individual GaN NR on graphene was contacted by placing one tungsten probe on top of the NR and another probe on graphene in its vicinity as seen in the SEM image in Figure 5b, at the left side. This configuration resulted in asymmetric I−V characteristics shown in Figure 5c (black curve). If the same procedure was applied to GaN NRs on the sapphire substrate in areas with no transferred graphene (the second probe is consequently placed on sapphire in the vicinity of the NR), no current could be measured (not shown in Figure 5c), demonstrating that the sapphire substrate remains insulating after the growth of GaN NRs. There is a difference in the slope for the contacted NR on graphene in the negative and positive voltage region. A similar nonlinear behavior is expected in a metal−semiconductor− metal model, where two Schottky contacts are standing opposite to each other with a semiconducting NR inbetween.40,41 If both Schottky contacts would have the same Schottky-barrier height, the resulting I−V characteristic would have a symmetrical slope in the positive and negative voltage regions. However, if the barrier heights at each Schottky contact are different, the I−V characteristics become asymmetric. The barrier heights can be influenced by different parameters like the work functions of the contacting materials, the contact area, and the doping level in the semiconductor. Due to the difference in the work function of tungsten (4.6 eV) and the electron affinity of GaN (4.1 eV), a Schottky contact is expected at this interface. The asymmetric I−V characteristic is indicative of a Schottky contact between the n-doped GaN NR and graphene as in the case of n-doped GaN layers.42 If both tungsten probes are placed on the same GaN NR (Figure 5b at the right side), the resulting I−V characteristic is symmetric again (Figure 5c, red curve), as it is expected to be for similar contact areas and work functions at both Schottky-contacts. In comparison to the contact via graphene, the current is considerably smaller (e.g., by over 2 orders of magnitude at

Figure 5. (a) Schematic drawing of the I−V measurement setup. The GaN NRs were contacted from above via a tungsten probe, while a second tungsten probe is positioned on graphene in the vicinity of the NRs. (b) SEM images of GaN NR for which the I−V characteristics were acquired. One nanoprobe on top of a GaN NR and the second nanoprobe on graphene (left). One nanoprobe on the top of a GaN NR and second near the base of the same NR (right). The scale bars in (b) represent 2 μm. (c) Semilogarithmic plot of the I−V characteristics measured for graphene and the GaN NR in (b).

−1 V). This can be the result of the smaller Schottky-barrier height between graphene and GaN as compared to GaN and the tungsten nanoprobes due to a larger contact area in the former case and differences in the work functions. These results show that GaN NRs grown on insulating substrates like sapphire can be contacted from underneath by inserting a graphene layer at the interface with the substrate before their actual growth. The tungsten nanoprobes were also used to remove single, as-grown GaN NRs from the substrates by bending them at the tip until they were breaking off, similar to experiments conducted on individual GaAs nanowires on graphite.18 In Figure 6, we present three representative cases, including one standing NR on sapphire (NR A) and two different NRs on graphene/sapphire (NR B and NR C). In the upper SEM images, the tungsten nanoprobe was in contact with the NRs without bending them yet. The central SEM images show the NRs before separating them from the substrate at their maximum deflection, while the lower SEM images display the respective areas after the fracture of the NRs. At the base of NR A, a small island of GaN material remained after the NR was F


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underlying substrate has a great influence on the regularity of the structures grown on graphene as evidenced by replacing sapphire with SiO2 and by studying the areas under the GaN NRs after their removal by tungsten nanoprobes. We propose that the NRs either grow inside nanoholes in graphene or on the top of graphene as implied by bending experiments. CL measurements demonstrated the good optical and structural properties of the upper part of NRs on graphene, with the defect luminescence being only detected near the lattice mismatched interface between the GaN NRs and graphene/ sapphire substrate. Raman spectroscopy proved the stability of graphene with respect to the high-temperature GaN growth, even beneath the GaN structures. Furthermore, it has been shown that GaN micro- and nanorods grow essentially on defective graphene, which provides the dangling bonds for the nucleation layer. The electric contact between the GaN NRs and graphene on the insulating sapphire substrate was confirmed through contacting the nanostructures from the top and graphene from the bottom via tungsten nanoprobes. These results could lead to a new kind of an atomically thin, electrical conducting mask for selective area growth by controlling the local defect configuration in pristine graphene. Thus, a spatially controlled nucleation and, consequently, growth of regular, electrically interconnected semiconductor nanorods for integration in future optoelectronic devices could be realized even on insulating substrates.

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AUTHOR INFORMATION

Corresponding Author

Figure 6. SEM images of three different NRs on sapphire (NR A) and on graphene (NR B and NR C) before bending (upper images), at maximum bending (central images), and after the breaking (lower images). The NRs were bent using a tungsten nanoprobe in a SEM. After the fracture, a GaN island remains attached to the sapphire substrate both without (NR A) and with graphene (NR B). It indicates that the growth of NR B occurred directly on the sapphire through nanoholes in graphene. Other NRs grew directly on graphene and were removed without leaving traces behind (NR C). The scale bars represent 2 μm.

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge the financial support from the German Research Foundation (DFG) within the research group “Dynamics and Interactions of Semiconductor Nanowires for Optoelectronics” (project number FOR 1616), the GRK 1896 “In-Situ Microscopy with Electrons, X-rays and Scanning Probes” as well as the “Excellence Cluster” “Engineering of Advanced Materials” at the FriedrichAlexander-Universität Erlangen-Nürnberg. The authors further acknowledge the European Union Seventh Framework Program (FP7/2007-2013) under the Grant n°280566, project UnivSEM.

removed, indicating the strong epitaxial bonding between the grown NRs and the sapphire substrate. After removing NR B from the graphene/sapphire substrate, such a remaining GaN spot can be seen as well, implying that such NRs nucleated on the sapphire substrate instead of graphene through nanoholes. If the NRs would nucleate on the graphene there should be no traces left behind due to the weaker bonding between GaN and graphene. Indeed, other NRs like NR C can be removed from graphene without leaving a visible GaN footprint, suggesting that such NRs grew on graphene, while still being vertically aligned, by means of van-der-Waals interactions with respect to the underlying graphene/sapphire substrate.43 Since no such NRs can be found on the graphene/SiO2 samples, these bending experiments show that the underlying substrate has a significant influence on the growth of micro- and nanostructures through/on the graphene covering the substrate. This important point is supported by Tsoi et al.44 that have shown that doped graphene is indeed transparent to the van der Waals forces of the underlying substrate.

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REFERENCES

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CONCLUSION In conclusion, we showed that vertically aligned and hexagonally shaped n-doped GaN NRs can be grown on single-layer graphene/sapphire substrates by MOVPE. The G


Article

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Growth of GaN Micro- and Nanorods on Graphene-Covered Sapphire

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