Vertically Oriented Growth of GaN Nanorods on Si ... - ACS Publications

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Vertically Oriented Growth of GaN Nanorods on Si Using Graphene as an Atomically Thin Buffer Layer Martin Heilmann,*,† A. Mazid Munshi,‡ George Sarau,†,§ Manuela Göbelt,† Christian Tessarek,†,§,∥ Vidar T. Fauske,⊥ Antonius T. J. van Helvoort,⊥ Jianfeng Yang,# Michael Latzel,†,∥ Björn Hoffmann,†,∥ Gavin Conibeer,# Helge Weman,‡,Δ and Silke Christiansen†,§,¶ †

Max Planck Institute for the Science of Light, Günther-Scharowsky-Str. 1, D-91058 Erlangen, Germany CrayoNano AS, Otto Nielsens vei 12, NO-7052 Trondheim, Norway § Institut für Nanoarchitekturen für die Energieumwandlung, Helmholtz-Zentrum Berlin für Materialien und Energie GmbH, Hahn-Meitner Platz 1, D-14109 Berlin, Germany ∥ Institute of Optics, Information and Photonics, Friedrich-Alexander-Universität Erlangen-Nürnberg (FAU), Staudtstr. 7/B2, D-91058 Erlangen, Germany ⊥ Department of Physics, Norwegian University of Science and Technology (NTNU), NO-7491 Trondheim, Norway # School of Photovoltaic and Renewable Energy Engineering, University of New South Wales, Kensington, Sydney, New South Wales 2052, Australia Δ Department of Electronics and Telecommunications, Norwegian University of Science and Technology (NTNU), NO-7491 Trondheim, Norway ¶ Physics Department, Freie Universität Berlin, Arnimallee 14, D-14195 Berlin, Germany ‡

S Supporting Information *

ABSTRACT: The monolithic integration of wurtzite GaN on Si via metal−organic vapor phase epitaxy is strongly hampered by lattice and thermal mismatch as well as meltback etching. This study presents single-layer graphene as an atomically thin buffer layer for c-axis-oriented growth of vertically aligned GaN nanorods mediated by nanometer-sized AlGaN nucleation islands. Nanostructures of similar morphology are demonstrated on graphene-covered Si(111) as well as Si(100). High crystal and optical quality of the nanorods are evidenced through scanning transmission electron microscopy, micro-Raman, and cathodoluminescence measurements supported by finite-difference time-domain simulations. Current−voltage characteristics revealed high vertical conduction of the as-grown GaN nanorods through the Si substrates. These findings are substantial to advance the integration of GaN-based devices on any substrates of choice that sustains the GaN growth temperatures, thereby permitting novel designs of GaN-based heterojunction device concepts. KEYWORDS: GaN, nanorods, graphene, MOVPE, GaN-on-Si

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several obstacles, such as high lattice mismatch, high thermal mismatch, plastic deformation, and meltback etching, making the growth of GaN on Si and vertical conduction challenging.3−5 A novel approach to counteract the lattice and thermal mismatch as well as the plastic deformation issues on Si, which usually result in high defect densities, is the fabrication of nanostructures such as nanorods (NRs), with large free surfaces being able to elastically relax the strain in all directions.6,7 GaN

ver the last two decades, GaN has evolved into one of the most important semiconductors after Si due to its outstanding optical and electrical properties, making it an attractive material for light-emitting diodes, lasers as well as high-power and high-frequency devices.1,2 Due to the lack of large native substrates, such GaN-based devices are usually grown heteroepitaxially on lattice mismatched substrates such as c-sapphire, 6H-SiC, or Si(111) using metal−organic vapor phase epitaxy (MOVPE) or molecular beam epitaxy (MBE). Among the different growth substrates, Si is most favorable due to its earth abundance, availability at low costs, large wafer diameters (up to 300 mm) and, most importantly, due to the possibility to monolithically integrate GaN with Si electronics. Using MOVPE, the most industrially viable method, there are © XXXX American Chemical Society

Received: February 3, 2016 Revised: April 18, 2016

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DOI: 10.1021/acs.nanolett.6b00484 Nano Lett. XXXX, XXX, XXX−XXX

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evidencing the NR’s high optical quality. The vertical conduction from the supporting Si(100) substrate throughout the GaN NRs are proven by contacting single, as-grown, ndoped GaN NRs from the top by a tungsten (W) nanoprobe. Commercially available single layer graphene on Si(111), Si(100), and c-sapphire was used as growth substrates. Prior to the growth of the NRs, a nitridation step is included in the recipes in which NH3 is supplied at high temperatures (details on the growth process can be found in the Method section). Density functional theory calculations predict chemical adsorption-like bonding for N atoms on graphene with a distance smaller than 2 Å, while for larger distances weaker physical adsorption-like bonding for adatoms is expected (e.g., for Al and Ga).26 To initiate the crystal growth on nitridated graphene, AlGaN nucleation islands are grown at 1200 °C at a gas phase ratio: trimethylaluminum (TMAl) and trimethylgallium (TMGa) of TMAl/(TMGa + TMAl) = 0.37. Using similar growth conditions for AlGaN thin film layers on graphene, Gupta et al.24 measured a solid Al composition of 29%. Al has a higher adsorption energy on graphene compared to Ga, which is beneficial for the adsorption via quasi van-der-Waals forces in the absence of defects in graphene.24,26 Moreover, Al shows a higher migration energy on graphene as compared to Ga, which supports island formation.26 The Al−N bond is stronger as compared to the Ga−N bond and desorption is less likely, making Al more favorable as well.27 Based on these points, we established a nucleation process that resides on the formation of AlGaN islands, which in turn support the growth of vertically aligned GaN NRs on graphene with high crystalline quality. The GaN NRs are grown at 1150 °C for 200 s (used for growth study on different substrate) and 400 s (employed for optical and electrical analyses). To obtain n-type doping, SiH4 was introduced into the reactor during GaN deposition. Recent reports demonstrate the crucial role of SiH4 not only for the controlled doping but also for achieving one-dimensionality of the GaN NR growth processes resulting in sidewall passivation of the growing NRs.28,29 Figure 1a and b show scanning electron microscopy (SEM) images of GaN NRs with identical morphology even though they were grown on differently orientated Si substrates covered with monolayer graphene. The hexagonally shaped NRs with varying height and diameter exhibit a pyramidal tip, indicating Ga-polarity. Not all AlGaN nucleation islands do initiate further GaN NR growth. Moreover, AlGaN islands nucleated in higher densities along lines in graphene (marked by orange arrows in Figure 1a and b). Such lines can result during the graphene growth process, where differently in-plane oriented graphene domains grow on the multicrystalline copper foil and extend until they coalesce.30 These grain boundaries are lines of higher defectiveness resulting in higher density of dangling bonds that can preferentially support the covalent bonding of atoms offered from the gaseous phase in the MOVPE processes. It was shown that AlN grown on defective step edges in epitaxial graphene tends to be multicrystalline.31 Thus, the growth of GaN is expected to be less favorable on grain boundaries in graphenecovered Si due to misalignment of the AlGaN lattice. On graphene-covered Si, the NRs grew in-between those grain boundaries with an average height and diameter of 300 ± 120 nm and 60 ± 30 nm, respectively. Their density of ∼2 × 108 cm−2 is considerably lower than the density of nucleation islands of ∼1.2 × 1010 cm−2. Besides the NRs, there are also large pyramidal shaped GaN microstructures with a density of ∼2 × 106 cm−2 (e.g., upper left corner of Figure 1a). A similar

NRs have been achieved by MOVPE and MBE on Si(111) using SiN masking layers or AlN buffer layers to prevent meltback etching, as a result of the reactions of Ga with Si when in intimate contact at high temperatures yielding local pit formation in Si.8−10 However, growth of wurtzite GaN on Si(100), and thus monolithically integration with already existing complementary metal−oxide−semiconductor (CMOS) technology, has been difficult due to the different crystal symmetries. In this respect, GaN nanowires can be grown on Si(100) using MBE, but nitridation of Si leads to the formation of an insulating SiN layer.11,12 Using a Si(100) wafer with a miscut of 4° in combination with a semi-insulating AlN seeding layer and a thick AlGaN buffer layer allowed for the MOVPE growth of GaN layers and their use in electronic devices.13,14 An intriguing approach to achieve high-quality GaN nanostructures on Si, or on any underlying substrate suitable for the epitaxy of GaN (e.g., sustaining high deposition temperatures up to ∼1200 °C), is the use of graphene as an intermediate buffer layer.15−18 The fabrication of large-area, high-quality graphene by chemical vapor deposition (CVD) and its transfer on various supporting substrates is achievable today with the prospective of graphene becoming a cheap and abundant substrate alternative.19,20 Graphene with its planar configuration of sp2-bonded carbon atoms can prevent meltback etching of Si21 and because of the hexagonal arrangement of C atoms, the one-atomic layer graphene itself can serve as a close to lattice matched substrate for the growth of wurtzite GaN.22 Growth of inclined and multifaceted GaN microstructures at temperatures below 1000 °C have been demonstrated by MOVPE on graphene-covered Si(111).17 Even despite the low quality of the crystal material, attributed to the low growth temperatures, a working photoconductive device could be realized.18 Thin, c-plane polycrystalline GaN layers grown by MBE were reported on graphene-covered Si(100).23 However, the GaN layers still suffered from stacking faults especially at the interface between graphene and GaN. In spite of the possible advantages of graphene as a growth substrate, the lack of dangling bonds in high-quality graphene leads to high surface tension and low nucleation densities. In this respect, it was already shown that GaN micro- and nanorods grown by MOVPE nucleate preferentially on defects in CVD graphene, which usually offer dangling bonds and support the GaN nucleation during epitaxy. An influence of the underlying, e.g., sapphire, substrate on the growth of GaN structures on top of the defective graphene was evident.16 AlN has shown to be a beneficial, yet semi-insulating, buffer layer on multilayered graphene, enabling covalent bonding for subsequent growth of III-nitride layers or dense nanostructures using MOVPE or MBE.15,24,25 In this work we demonstrate a MOVPE growth process, initiated by nanometer-sized AlGaN nucleation islands, for caxis-oriented, single crystalline GaN NRs that grow vertically on single-layer graphene on Si(111), Si(100), and c-sapphire. Single crystallinity and [0001] orientation of GaN NRs are demonstrated by cross-sectional scanning transmission electron microscopy (STEM). Micro-Raman spectroscopy investigations along the growth axis of separated nanorods indicate an increasing level of residual strain toward the interface. Furthermore, they confirm the high crystalline quality of the single-layer graphene before and after the growth of the GaN NRs. Cathodoluminescence (CL) measurements of individual NRs combined with finite-difference time-domain (FDTD) simulations show the existence of fundamental cavity modes, B


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to the underlying sapphire substrate (Figure S1). On the bare Si(111) reference substrate no NRs grew under similar growth conditions (Figure 1d). Instead, misaligned microstructures as well as meltback etching are observed. The findings demonstrate that graphene as an atomically thin buffer layer can prevent alloying between Ga and Si during the MOVPE growth of GaN NRs on Si. Based on the aforementioned observations, it is noteworthy that AlGaN nucleation islands on graphene enabled the vertically aligned, single crystalline growth of GaN NRs on different supporting substrates. This clearly differentiates the here proposed growth procedure for GaN NRs on graphene from previous published literature, where the NRs were either misaligned17,18,34 or growing on defects in graphene with an influence of the supporting substrate.16 Since Si(100) is widely used as industrystandard substrate in Si-based technologies and growth of high quality wurtzite GaN on it is more challenging, we will focus in the following on GaN NRs grown on graphene-covered Si(100) substrates. Cross-sectional high-resolution STEM (HRSTEM) measurements revealed the single crystallinity of the GaN NRs on graphene-covered Si(100), which are predominantly free of extended structural defects with no threading dislocations, inversion domain boundaries, or stacking faults visible in the upper part of the NRs (see Figure 2a).23,32,35 In the bright-field (BF) STEM image of the interface (Figure 2b) one can clearly distinguish between the Si substrate, the single layer graphene, and the pristine GaN NR. Moreover, an amorphous layer (∼2 nm) between the crystalline Si substrate and graphene is visible. This amorphous layer can be attributed to the native silicon oxide (SiOx). Investigating the same interface region under high-angle annular dark field (HAADF) conditions showed an area of lower intensity in the center of the NR (Figure 2c). This is attributed to strain and lattice distortion of the GaN at the interface to the AlGaN as well as the atomic number difference between Al and Ga.36 Thus, the AlGaN nucleation islands appear slightly darker than the surrounding GaN material. Superimposing annular bright-field (ABF) STEM and annular dark-field (ADF) STEM images of the same sample area permits a direct determination of the polarity of the NR. ADF images are more sensitive to heavy elements, revealing the position of Ga with respect to lighter N when superimposed on the ABF image, where both Ga and N are visible.37 Figure 2d, acquired from the upper part of a NR, shows the ABF image in combination with a Ga-polar GaN model, proving the Gapolarity of the grown GaN NRs on graphene. Figure 2e displays the HAADF STEM image of the yellow framed area in Figure 2c for which the energy dispersive X-ray (EDX) Kα-signals of Al, Si, and O were superimposed. The respective normalized Kα-signal intensities of these elements as well as Ga and N in the region limited by the dashed black lines in e are shown in Figure 2f. A higher O signal from the interface below graphene was detected but no Al, N, or Ga. These EDX results prove that the GaN NR grew over the AlGaN islands and confirm the amorphous layer below graphene to be SiOx. Graphene as diffusion barrier can prevent O to be removed from the SiOx layer on the Si surface during the initial bakeout process. Furthermore, graphene can block diffusion of Ga and N into the Si and hence circumvent meltback etching of the Si substrate and the formation of a SiN layer. For the investigation of the strain distribution along the growth axis within a single NR using nonresonant microRaman spectroscopy, the longer NRs were studied. The

Figure 1. SEM images of GaN NRs grown on AlGaN nucleation islands on different substrates: (a) graphene-covered Si(111), (b) graphene-covered Si(100), (c) graphene-covered c-sapphire, and (d) Si(111) without graphene where meltback etching is observed. Orange arrows in a and b indicate the positions of wrinkles or coalescence lines of different oriented domains in graphene.

formation of N-polar GaN microstructures, evolving during self-catalyzed growth of NRs on sapphire has been observed.32 The influence of the underlying substrate on the epitaxy of GaN NRs on top of graphene was studied by applying the same growth conditions with AlGaN nucleation islands on graphenecovered sapphire (see Figure 1c). Similar to a previous study, predominantly N-polar GaN micro- and nanostructures with flat top facets grew on lines of darker SEM contrast attributed to highly defective graphene on insulating sapphire.16 It seems that the underlying substrate can influence the GaN epitaxy on graphene at defective lines, possibly through nanoholes in graphene.16 Such nanoholes can allow growth of N-polar NRs on bare, nitridated sapphire, which seems not possible at nanoholes in graphene-covered Si. Nevertheless, besides the Npolar GaN NRs one can see NRs with a pyramidal top facet as well as pyramidal microstructures, implying Ga-polar growth in areas of bright SEM contrast corresponding to graphene with very low defect density, similar to NRs grown on graphenecovered Si. Ga-polar NRs are smaller as compared to N-polar NRs, regarding the height and diameter. A difference in decomposition and sublimation between Ga- and N-polar GaN at elevated temperatures in an H2 and NH3 atmosphere can be responsible for the different sizes of the structures.33 Moreover, no preferential in-plane orientation of the Ga-polar NRs was observed, while the N-polar NRs at defective lines on graphenecovered sapphire are oriented all in the same direction. This alignment is most likely mediated by the epitaxial relationship C


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Figure 2. (a) Cross-sectional HRSTEM image of a GaN NR on graphene-covered Si(100). (b, c) Simultaneously acquired BF and HAADF STEM images of the interface between GaN, graphene, and Si (blue dashed region in a). (d) High-magnification ABF STEM image from the upper part of the NR superimposed with a model for Ga-polar GaN. (e) HAADF STEM image of the interface (yellow dashed region in c) superimposed with the EDX signal of Al, O, and Si. (f) Integrated linescan of the Kα-signal of Ga, Al, N, O, and Si from the area between the vertical black dashed lines in d showing no Ga, N, or Al signal below the graphene.

growth and the AlGaN nucleation island. At the interface between AlGaN and GaN thermal expansion coefficient mismatch leads to tensile strain and lattice distortion in our structures. Experimentally, a Raman shift of −2.57 ± 0.05 cm−1 was found in GaN layers grown on Si(111) substrates using a 5 nm thin AlN buffer layer, which can be translated as an in-plane biaxial tensile strain of 0.22%, being attributed to lattice and thermal mismatch between these materials.39,41 Even though, based on the lattice constants of GaN and AlGaN, one would expect a compressive strain in GaN, the experimentally measured tensile strain emphasizes the complex nature of the three-dimensional strained state at such an interface. Moreover, there are varying values for the thermal expansion coefficient of graphene itself42 and an additional influence of the underlying substrate on this coefficient is still not conclusively answered.43 A standard approach to reduce the strain in GaN thin films on Si is to increase the thickness of the complex AlN and AlGaN buffer layered structures.5 Nevertheless, using graphene as a buffer layer on Si(100) we measured an even smaller Raman shift of −0.76 ± 0.08 cm−1, equivalent to as low residual tensile strain as 0.067% at the base of the GaN NRs grown on nanometer-sized AlGaN nucleation islands. Raman mapping measurements were further utilized for the characterization of graphene before and after the GaN NRs growth. It should be noted that the Raman signal of graphene was detected from every probed spot, even though graphene was not visible besides the NRs in the STEM images (see Figure 2b). The mean spectra in Figure 3c show the characteristic D, G, and 2D peaks of graphene before/after

increased growth time from 200 to 400 s at otherwise identical growth conditions resulted in increased lengths and diameters of the NRs at the expense of their density, with average values of: length = 1650 ± 550 nm, diameter = 150 ± 60 nm, and density ∼2 × 107 cm−2. A similar behavior of reduction of NR density with growth time was reported for self-catalyzed growth of GaN NRs on sapphire38 and explained by Ostwald ripening of the Ga droplet and dissolution of NRs without Ga droplet at their tips during the vapor−liquid−solid growth at elevated growth temperatures.32 Raman spectra were acquired stepwise over the lengths of well-separated and substrate attached NRs located at the sample edge employing a 90° tilted holder. Figure 3a shows a Raman spectrum acquired on the top of a representative NR along with a schematic of the side measurement geometry in the inset. Besides the strain-sensitive E2h peak at 567.68 cm−1 corresponding to strain-free GaN bulk material, one can also identify the E1(TO) peak at 557.67 cm−1, the quasi-TO mode at 548.13 cm−1, and the A1(TO) peak at 530.09 cm−1.39 An additional, peak belonging to strain-free, crystalline Si is visible at 521.20 cm−1 indicating that the GaN NR can act as a waveguide for the used laser excitation light (λ = 457 nm) in the micro-Raman measurements when focusing on the NR top in this geometry. While approaching the bottom of the NRs, the E2h peak positions shift toward lower wavenumbers consistent with the presence of residual tensile strain, as displayed in Figure 3b (the data points represent the average values over five NRs of similar length).40 According to the STEM measurements in Figure 2, the GaN NRs are in contact with the single layer graphene as a result of the lateral D


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conditions on Si. It was shown that graphene resists the GaN NR growth conditions when being transferred to sapphire substrates.16 However, the thermal expansion of sapphire and Si are markedly different, and thus these results demonstrate that graphene can act as a buffer layer on substrates with different thermal expansion coefficients. The optical properties of single, as-grown GaN NRs were characterized using room temperature CL, and the results obtained for one representative NR are summarized in Figure 4. The NR depicted in Figure 4a with a height of 2.1 μm and a diameter of 200 nm shows a bright CL signal in the panchromatic CL mapping, which is getting weaker toward the bottom of the NR (see Figure 4b). The monochromatic CL mapping of the near band edge emission (NBE) at 365 nm (Figure 4c) shows a lower intensity toward the bottom as well, while the yellow defect luminescence (YL) at 560 nm has a homogeneous intensity distribution along the NR (Figure 4d). Point defects, especially at the surface of grown nanostructures or defects introduced through Si doping, can be responsible for the YL.48−50 Differently colored arrows in the SEM image indicate the positions at which the electron beam was focused during spectral acquisition. The semilogarithmic plot of the spectrum from the whole NR (depicted by the dashed pink box in Figure 4a) displays a broad NBE around 3.4 eV and a relatively weak YL around 2.2 eV (Figure 4e, box) superimposed with additional peaks. The NBE emission exhibits a shift (figuratively depicted through the dashed red line in Figure 4e) to higher energies (∼40 meV) as well as a broadening of the fwhm (∼110 meV) from the bottom toward the top of the NR. A shift of the NBE can be explained by the Burstein−Moss effect, which was also observed in GaN microrods grown on sapphire with different levels of Si-supply during growth.51,52 For countable blue shifting of the NBE, the electron concentration is expected to exceed 9 × 1018 cm−3 at which the Burstein−Moss effect is supposed to overcome the band edge renormalization at very high free electron densities.51 Additionally, the relaxation of tensile strain from the bottom to the top of the NRs, as evidenced in the Raman measurements (Figure 3b), can contribute to the shift of the NBE to higher energies.53 A broadening of the fwhm of the NBE for increasing Si doping concentrations in GaN was observed at room temperature by Schubert et al.54 and attributed to potential fluctuations caused by an inhomogeneous distribution of the Si dopants. Since the monochromatic CL mapping of the YL showed a homogeneous intensity distribution along the NR (Figure 4d) and no extended defects were found in the upper part of the NRs (Figure 2a), the increasing broadening of the NBE from bottom to top is attributed to an increasing doping concentration. Additional peaks superimposed to the NBE affect a clear determination of the fwhm. FDTD simulations were conducted for a NR with similar geometry to study the nature of the superimposed peaks. A dipole source was placed at the top of the NR and the far-field emission spectra are calculated by integrating the Poynting vector over the upper half space, a comparable situation to the CL measurements. A constant refractive index of n = 2.5 for GaN was assumed for all wavelengths and the in-plane (xy) and longitudinal polarized (z-direction along NR c-axis) emissions of the source are separately considered. The electric field intensity profile in the xz-plane of the NR (Figure 4f) exhibits an increasing number of nodes along the c-axis for higher energies (from 15 to 20 nodes in the range from 2.884 to 3.263 eV) showing that the NR can act as an optical cavity for the

Figure 3. (a) Raman spectrum acquired from the top of the GaN NR displayed in the SEM image in b with the characteristic peaks from GaN as well as from Si. (b) Following the position of the E2h-peak permits a strain estimation along the NRs c-axis, showing an increased strain toward the bottom of the NR, near the interface to graphene. Scale bar in the SEM image is 1 μm. (c) Mean Raman spectra of graphene before and after the GaN NR growth, showing D, G, and 2D peaks characteristic for single layer graphene (the spectra are vertically shifted for clarity). Insets in a and c depict the Raman measurement configurations.

GaN NR growth with fitting values for the 2D to G ratio of ∼2.4/1.58, the 2D peak full width at half-maximum (fwhm) of ∼34/42 cm−1, and the D to G ratio of ∼0.05/0.07. These parameters imply that a high-quality single layer graphene with very low defect density is preserved even after the high temperature exposure to the MOVPE precursor gases (including the N-doping of graphene).44−47 Thus, the Raman measurements not only confirm that GaN NRs grow nearly relaxed with only a mild tensile strain on the graphene-covered Si, but also that graphene itself is resilient toward the growth E


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Figure 4. (a) SEM image of the investigated GaN NR. Arrows indicate the fixed positions of the electron beam during the acquisition of the CL spectra. (b) Correlated panchromatic CL intensity map of the same sample area. (c, d) Monochromatic CL intensity map acquired at 365 and 560 nm, respectively. (e) Normalized semilogarithmic plot of CL spectra of the whole NR (box) and for a fixed electron beam at different positions along the NR (top, center, and bottom). Figuratively the red dashed line depicts the shift of the NBE. (f) Comparison of electric field intensity profiles in the xz-plane simulated using FDTD for the energies denoted in (e) top as gray values. (g) Electric field intensity profile of the HE11 mode in the xyplane with the dipole being symmetric to the y- (left) and x-axis (right). The white dashed lines depict the boundaries of the NR.

ments. The contact between the stage and the Si substrate was realized through Ga/In eutectic. Figure 5a (left) shows an SEM image of a GaN NR shortly before it is in contact with the W nanoprobe. Since W with a work function of 5.1 eV is unlikely to form an Ohmic contact to GaN with an electron affinity of 4.1 eV, a Schottky-contact is expected. To the right of Figure 5a current-density (J) vs voltage (V) characteristics of several NRs with different slopes are displayed. The inset shows a schematic of the measurement configuration. For the studied NRs with diameters around 400 nm and length of around 1.2 μm, current-densities of up to 14 kA/cm2 were measured at 3 V. The native silicon oxide detected by TEM (Figure 2) beneath the graphene should substantially influence the JV characteristics in this configuration. Therefore, to demonstrate vertical conduction, another sample was prepared for which the native oxide on degenerately p-doped Si (5 × 1018 cm−3) was removed by a HF etch (see Methods) prior to the transfer of graphene. The same growth conditions with a GaN NR growth time of 400 s were applied. The JV characteristics of such NRs

respective energies (corresponding to the experimental CL peaks denoted in gray in Figure 4e top). The profile of the fundamental HE11 waveguide mode in the xy-plane is shown in Figure 4g for two different orientations of the dipole, either symmetric to the y- or x-axis, respectively. Fundamental cavity modes along the NR, as identified through the FDTD simulation, support the assumption of waveguiding within the NR, implied from the micro-Raman measurements, and confirm the high crystal quality of the NRs grown on graphene. The use of GaN based devices grown on insulating sapphire usually requires either double front contacts or a flip-chip approach in which the active GaN structures are removed from the substrate. Even if the GaN structure is grown on a conducting Si substrate, a thick buffer layer, containing threading dislocations, impedes vertical conduction.5 For electrical characterization single, as-grown GaN NRs were contacted from the top via a W nanoprobe and from the back via the Si substrate inside an SEM chamber at room temperature. The electron beam was off during the measureF


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Figure 5. (a) SEM image of a GaN NR and the W nanoprobe above it (left). Nonsymmetric current density vs voltage characteristics of several NRs on graphene-covered, degenerately n-doped Si(100) with native oxide (right). Inset: Schematic of the measurement configuration. (b) SEM image of a GaN NR on degenerately p-doped Si(100), which was HF-etched prior to graphene transfer and the W nanoprobe above it (left). Current density vs voltage plot (blue circles) for a contacted NR (center). The solid orange line is a fit using a MSM model. Inset: Schematic of the measurement configuration. At current densities above 400 kA/cm2 resistive heating led to decomposition of the GaN NR material itself directly at the Schottkyjunction to the nanoprobe and on the other hand to the formation of Ga droplets (right).

also serves as an efficient chemically inert buffer, preventing the diffusion of Ga into the Si and thus circumventing meltback etching of Si substrates, which is a severe problem in MOVPE growth of GaN on Si. The nearly strain-free vertical GaN NRs were shown to support fundamental optical cavity modes, evidencing their good optical properties. Vertical conduction was proven through contacting individual, as-grown n-doped GaN NRs with a W nanoprobe from the top and via the degenerately p-type doped Si substrate from the back. The finding that graphene can effectively act as an epitaxy mediating layer for subsequent growth of GaN NRs opens up a wealth of options to basically grow such nanostructures on any hightemperature sustaining growth substrate, including cheap and large area substrates that merely fulfill the role of being a mechanical support. Methods. Substrate Preparation and GaN NR Growth. Commercially available single layer CVD graphene transferred on c-sapphire, Si(111), and Si(100) substrates was used for the growth of GaN NRs by MOVPE (Aixtron 200RF horizontal flow reactor). The growth temperature was monitored using a thermocouple inside the susceptor with the substrates surface temperature on top of it being approximately 50 to 100 °C lower. All growth steps were conducted at 100 mbar. Preceding the growth a cleaning step was conducted in H2 atmosphere at 1200 °C for 5 min followed by a nitridation step using ammonia (NH3) with a flow of 600 sccm for 10 min. Subsequently AlGaN nucleation islands were grown on graphene for 40 s by introducing trimethylgallium (TMGa)

with a similar geometry (SEM image on the left-hand side of Figure 5b) as in the previous sample show a nonlinearity as well as a significantly enhanced current-density of 350 kA/cm2 at 3 V (Figure 5b center). Similarly, nonlinear current−voltage (IV) curves were observed for GaAs NRs grown via MBE on highly doped Si and were attributed to Schottky junctions at both ends of the NRs.55,56 A metal−semiconductor−metal (MSM) model can be used to determine the charge-carrier density and the carrier mobility for this configuration (a description of the model can be found in the Supporting Information).57−59 Here we applied a fitting routine based on the MSM model developed by Liu et al.58 by considering the W nanoprobe and the graphene-covered Si substrate as Schottky junctions to the GaN NRs and retrieved charge-carrier densities and mobilities of 1.5 ± 0.2 × 1018 cm−3 and 108 ± 12 cm2/(V s), respectively. A conductivity of 26.80 ± 0.40 S/cm derived from the fit directly proves the vertical conduction of the asgrown GaN NRs. However, resistive heating at the Schottky junction between GaN and the W nanoprobe led to decomposition of the NRs at a forward bias above 3.5 V (see SEM image on the right-hand side of Figure 5b) and thus limited measurements at higher voltages. In conclusion, vertical growth of GaN NRs by MOVPE on CVD graphene could be successfully obtained as mediated by nanometer-sized AlGaN nucleation islands. A high quality integration of GaN on Si(100) is shown to be viable using the atomically thin graphene interlayer that serves as a buffer layer for GaN growth. In addition to serving as a substrate, graphene G


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Nano Letters and trimethylaluminum (TMAl) with a flux of 44.8 and 26.3 μmol/min, respectively. The NH3 flow was set to 600 sccm and kept constant after the growth of the nucleation islands for stabilizing them while the temperature was lowered to 1150 °C. GaN NRs were grown using the same TMGa flow as before but with only 25 sccm NH3 for 200 s. Silane was offered during the NR growth for n-type doping and sidewall passivation with an effective flow of 0.03 μmol/min. While cooling down the MOVPE reactor, the NH3-flux was kept constant at 25 sccm. For the optical characterization longer NR were grown for 400 s under the same growth conditions. For electrical characterization additional samples were grown for 400 s on Cu catalyzed CVD graphene transferred on Si(100) substrates, which were etched in HF (5%) for 2 min prior to the transfer. A description of the growth parameters of the Cu-catalyzed CVD graphene as well as the transfer process can be found elsewhere.16 Scanning Electron Microscopy and Cathodoluminescence Measurements. SEM images were acquired using a Hitachi S4800 at an acceleration voltage of 5 keV at an angle between surface normal and electron beam of 60°. For correlated CL and SEM measurements, a Tescan MIRA3 GM at an acceleration voltage of 5 keV was used. The angle between the surface normal and electron beam was adjusted to 70°. The SEM was equipped with a CL unit (Gatan MonoCL), and the measurements were conducted at room temperature. The integration time for the spectral acquisition was 40 s, and a monochromator grating with 300 l/mm was used. Transmission Electron Microscopy. TEM-lamella were prepared with a Ga2+-ion beam using a FEI Helios DualBeam at 5−30 keV. The structures where protected through C and Pt protection layers. The TEM analysis was done in a double corrected JEOL ARM 200F TEM/STEM at 80 keV operated in STEM mode with a 100 mm2 JEOL Centurio EDX SDD. The EDX data were analyzed using HyperSpy.60 Raman Spectroscopy. Nonresonant micro-Raman measurements were carried out at room temperature in backscattered configuration using a LabRam HR800 (Horiba Scientific). A linear polarized laser with a wavelength of 457 nm was employed for the Raman excitation. The laser light was focused by a 100× objective (numerical aperture 0.9), resulting in a spot size of ∼700 nm corresponding to a laser power on the sample surface of ∼600 μW. No structural changes to the graphene and the GaN NRs as well as no Raman shift due to local heating by the laser beam were observed. Nanoprobing. Single GaN NRs and the Si substrate were contacted via a W nanoprobe from the top and via Ga/In contact from the back. The samples were tilted by 60° between the surface normal and the electron beam. The nanoprobe was maneuvered inside the SEM (FEI Strata DB 235) using an integrated-circuit testing device with motorized probing arms (Kammrath & Weiss). Current−voltage characteristics were recorded at room temperature with a semiconductor characterization system (Keithley SCS 4200). Finite-Difference Time-Domain Simulations. Lumerical FDTD Solutions61 was used to identify the optical modes supported in the NR structure. The size and morphology of the simulated NR were retrieved through SEM measurements. An electric dipole source was used to simulate the re-emission process after the excitation with the electron beam at the top of the NR. For simplicity the refractive index of the NR was set to n = 2.5 and to n = 1 for the surrounding for all wavelengths. A power monitor was placed around the upper hemisphere, with a

distance of one wavelength, to integrate the Poynting vector and acquire the far-field emission spectra. The in-plane polarized (x- and y-direction) and longitudinal polarized (zdirection) emission are separately considered.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.nanolett.6b00484. SEM images to determine the in-plane orientation of GaN NRs on graphene-covered substrates. Description of the MSM model. (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected], martinheilmann@gmx. de. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS M.H., G.S., M.G., C.T., M.L., B.H., and S.C. gratefully acknowledge the financial support from the German Research Foundation (DFG) within the research projects “Dynamics and Interactions of Semiconductor Nanowires for Optoelectronics” (FOR 1616), “Hybrid Inorganic/Organic Systems for Optoelectronics” (HIOS, SFB 951), “In-Situ Microscopy with Electrons, X-rays and Scanning Probes” (GRK 1896) as well as the cluster of excellence “Engineering of Advanced Materials” at the Friedrich-Alexander-Universität ErlangenNürnberg. We would also like to thank the “Deutscher Akademischer Austauschdienst” (DAAD) for providing exchange scholarship. The authors further acknowledge the European Union Seventh Framework Program (FP7/20072013) under the grant agreement no. 280566, project UnivSEM. V.T.F., A.T.J.v.H., and H.W. acknowledge the support from the Research Council of Norway (RCN) through the FRINATEK (Grant 214235) and NANO2021 (Grant 239206) programs, as well as for the support to NTNU NanoLab through the Norwegian Micro- and Nano-Fabrication Facility, NorFab (Grant 197411/V30), and the NORTEM project (Grant 197405).

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REFERENCES

(1) Nakamura, S.; Fasol, G. The blue laser diode: GaN based light emitters and lasers; Springer: Berlin, 1997. (2) Quay, R. Gallium nitride electronics; Springer: Berlin, 2008. (3) Dadgar, A.; Krost, A. In III-Nitride Semiconductors and their Modern Devices; Gil, B., Ed.; Oxford University Press: Oxford, 2013. (4) Dadgar, A. Phys. Status Solidi B 2015, 252, 1063−1068. (5) Zhu, D.; Wallis, D. J.; Humphreys, C. J. Rep. Prog. Phys. 2013, 76, 106501. (6) Mårtensson, T.; Svensson, C. P. T.; Wacaser, B. A.; Larsson, M. W.; Seifert, W.; Deppert, K.; Gustafsson, A.; Wallenberg, L. R.; Samuelson, L. Nano Lett. 2004, 4, 1987−1990. (7) Kavanagh, K. L. Semicond. Sci. Technol. 2010, 25, 24006. (8) Salomon, D.; Dussaigne, A.; Lafossas, M.; Durand, C.; Bougerol, C.; Ferret, P.; Eymery, J. Nanoscale Res. Lett. 2013, 8, 61. (9) Bertness, K.; Sanford, N.; Barker, J.; Schlager, J.; Roshko, A.; Davydov, A.; Levin, I. J. Electron. Mater. 2006, 35, 576−580.

H


Letter

Nano Letters (10) Koester, R.; Sager, D.; Quitsch, W.; Pfingsten, O.; Poloczek, A.; Blumenthal, S.; Keller, G.; Prost, W.; Bacher, G.; Tegude, F. Nano Lett. 2015, 15, 2318−2323. (11) Cerutti, L.; Ristić, J.; Fernández-Garrido, S.; Calleja, E.; Trampert, A.; Ploog, K. H.; Lazic, S.; Calleja, J. M. Appl. Phys. Lett. 2006, 88, 160−163. (12) Chèze, C.; Geelhaar, L.; Trampert, A.; Riechert, H. Appl. Phys. Lett. 2010, 97, 043101. (13) Schulze, F.; Dadgar, A.; Bläsing, J.; Hempel, T.; Diez, A.; Christen, J.; Krost, A. J. Cryst. Growth 2006, 289, 485−488. (14) Schulze, F.; Kisel, O.; Dadgar, A.; Krtschil, A.; Bläsing, J.; Kunze, M.; Daumiller, I.; Hempel, T.; Diez, A.; Clos, R.; Christen, J.; Krost, A. J. Cryst. Growth 2007, 299, 399−403. (15) Hayashi, H.; Konno, Y.; Kishino, K. Nanotechnology 2016, 27, 055302. (16) Heilmann, M.; Sarau, G.; Göbelt, M.; Latzel, M.; Sadhujan, S.; Tessarek, C.; Christiansen, S. Cryst. Growth Des. 2015, 15, 2079−2086. (17) Kang, S.; Mandal, A.; Park, J.-H.; Um, D.-Y.; Chu, J. H.; Kwon, S.-Y.; Lee, C.-R. J. Alloys Compd. 2015, 644, 808−813. (18) Kang, S.; Mandal, A.; Chu, J. H.; Park, J.-H.; Kwon, S.-Y.; Lee, C.-R. Sci. Rep. 2015, 5, 10808. (19) Li, X.; Cai, W.; An, J.; Kim, S.; Nah, J.; Yang, D.; Piner, R.; Velamakanni, A.; Jung, I.; Tutuc, E.; Banerjee, S. K.; Colombo, L.; Ruoff, R. S. Science 2009, 324, 1312−1314. (20) Bae, S.; Kim, H.; Lee, Y.; Xu, X.; Park, J.; Zheng, Y.; Balakrishnan, J.; Lei, T.; Kim, H. R.; Song, Y. Il; Kim, Y.; Kim, K. S. Nat. Nanotechnol. 2010, 5, 1−5. (21) Chen, S.; Brown, L.; Levendorf, M.; Cai, W.; Ju, S.; Edgeworth, J.; Li, X.; Magnuson, C. W.; Velamakanni, A.; Piner, R. D.; Kang, J.; Park, J.; Ruoff, R. S. ACS Nano 2011, 5, 1321−1327. (22) Munshi, A. M.; Dheeraj, D. L.; Fauske, V. T.; Kim, D.-C.; van Helvoort, A. T. J.; Fimland, B.-O.; Weman, H. Nano Lett. 2012, 12, 4570−4576. (23) Araki, T.; Uchimura, S.; Sakaguchi, J.; Nanishi, Y.; Fujishima, T.; Hsu, A.; Kim, K. K.; Palacios, T.; Pesquera, A.; Centeno, A.; Zurutuza, A. Appl. Phys. Express 2014, 7, 071001. (24) Gupta, P.; Rahman, A. A.; Hatui, N.; Gokhale, M. R.; Deshmukh, M. M.; Bhattacharya, A. J. Cryst. Growth 2013, 372, 105−108. (25) Kovács, A.; Duchamp, M.; Dunin-Borkowski, R. E.; Yakimova, R.; Neumann, P. L.; Behmenburg, H.; Foltynski, B.; Giesen, C.; Heuken, M.; Pécz, B. Adv. Mater. Interfaces 2015, 2, 1400230. (26) Nakada, K.; Ishii, A. Solid State Commun. 2011, 151, 13−16. (27) Iliopoulos, E.; Moustakas, T. D. Appl. Phys. Lett. 2002, 81, 295− 297. (28) Tessarek, C.; Heilmann, M.; Butzen, E.; Haab, A.; Hardtdegen, H.; Dieker, C.; Spiecker, E.; Christiansen, S. Cryst. Growth Des. 2014, 14, 1486−1492. (29) Hartmann, J.; Wang, X.; Schuhmann, H.; Dziony, W.; Caccamo, L.; Ledig, J.; Mohajerani, M. S.; Schimpke, T.; Bähr, M.; Lilienkamp, G.; Daum, W.; Seibt, M.; Straßburg, M.; Wehmann, H.-H.; Waag, A. Phys. Status Solidi A 2015, 212, 2830−2836. (30) Kidambi, P. R.; Bayer, B. C.; Blume, R.; Wang, Z.-J.; Baehtz, C.; Weatherup, R. S.; Willinger, M.-G.; Schloegl, R.; Hofmann, S. Nano Lett. 2013, 13, 4769−4778. (31) Al Balushi, Z. Y.; Miyagi, T.; Lin, Y.-C.; Wang, K.; Calderin, L.; Bhimanapati, G.; Redwing, J. M.; Robinson, J. A. Surf. Sci. 2015, 634, 81−88. (32) Tessarek, C.; Bashouti, M.; Heilmann, M.; Dieker, C.; Knoke, I.; Spiecker, E.; Christiansen, S. J. Appl. Phys. 2013, 114, 144304. (33) Mastro, M. A.; Kryliouk, O. M.; Anderson, T. J.; Davydov, A.; Shapiro, A. J. Cryst. Growth 2005, 274, 38−46. (34) Chung, K.; Beak, H.; Tchoe, Y.; Oh, H.; Yoo, H.; Kim, M.; Yi, G.-C. APL Mater. 2014, 2, 092512. (35) Bergbauer, W.; Strassburg, M.; Kölper, C.; Linder, N.; Roder, C.; Lähnemann, J.; Trampert, A.; Fündling, S.; Li, S. F.; Wehmann, H.H.; Waag, A. Nanotechnology 2010, 21, 305201. (36) Shiojiri, M.; Č eh, M.; Šturm, S.; Chuo, C. C.; Hsu, J. T.; Yang, J. R.; Saijo, H. J. Appl. Phys. 2006, 100, 013110.

(37) De la Mata, M.; Magen, C.; Gazquez, J.; Utama, M. I. B.; Heiss, M.; Lopatin, S.; Furtmayr, F.; Fernández-Rojas, C. J.; Peng, B.; Morante, J. R.; Rurali, R.; Eickhoff, M.; Fontcuberta i Morral, A.; Xiong, Q.; Arbiol, J. Nano Lett. 2012, 12, 2579−2586. (38) Tessarek, C.; Dieker, C.; Spiecker, E.; Christiansen, S. Jpn. J. Appl. Phys. 2013, 52, 08JE09. (39) Sarau, G.; Heilmann, M.; Latzel, M.; Christiansen, S. Nanoscale 2014, 6, 11953−11962. (40) Sarau, G.; Christiansen, S.; Holla, M.; Seifert, W. Sol. Energy Mater. Sol. Cells 2011, 95, 2264−2271. (41) Bairamov, B. H.; Gürdal, O.; Botchkarev, A.; Morkoc, H.; Irmer, G.; Monecke, J. Phys. Rev. B: Condens. Matter Mater. Phys. 1999, 60, 16741−16746. (42) Magnin, Y.; Förster, G. D.; Rabilloud, F.; Calvo, F.; Zappelli, a; Bichara, C. J. Phys.: Condens. Matter 2014, 26, 185401. (43) Linas, S.; Magnin, Y.; Poinsot, B.; Boisron, O.; Martinez, V.; Fulcrand, R.; Tournus, F.; Dupuis, V.; Rabilloud, F.; Bardotti, L.; Han, Z.; Kalita, D.; Bouchiat, V.; Calvo, F. Phys. Rev. B: Condens. Matter Mater. Phys. 2015, 91, 1−5. (44) Das, A.; Pisana, S.; Chakraborty, B.; Piscanec, S.; Saha, S. K.; Waghmare, U. V.; Novoselov, K. S.; Krishnamurthy, H. R.; Geim, A. K.; Ferrari, A. C.; Sood, A. K. Nat. Nanotechnol. 2008, 3, 210−215. (45) Wang, Q. H.; Jin, Z.; Kim, K. K.; Hilmer, A. J.; Paulus, G. L. C.; Shih, C.-J.; Ham, M.-H.; Sanchez-Yamagishi, J. D.; Watanabe, K.; Taniguchi, T.; Kong, J.; Jarillo-Herrero, P.; Strano, M. S. Nat. Chem. 2012, 4, 724−732. (46) Zafar, Z.; Ni, Z. H.; Wu, X.; Shi, Z. X.; Nan, H. Y.; Bai, J.; Sun, L. T. Carbon 2013, 61, 57−62. (47) Sarau, G.; Lahiri, B.; Banzer, P.; Gupta, P.; Bhattacharya, A.; Vollmer, F.; Christiansen, S. Adv. Opt. Mater. 2013, 1, 151−157. (48) Reshchikov, M. A.; Morkoç, H. J. Appl. Phys. 2005, 97, 061301. (49) Huang, P.; Zong, H.; Shi, J.-J.; Zhang, M.; Jiang, X.-H.; Zhong, H.-X.; Ding, Y.-M.; He, Y.-P.; Lu, J.; Hu, X.-D. ACS Nano 2015, 9, 9276−9283. (50) Li, Q.; Wang, G. T. Nano Lett. 2010, 10, 1554−1558. (51) Feneberg, M.; Osterburg, S.; Lange, K.; Lidig, C.; Garke, B.; Goldhahn, R.; Richter, E.; Netzel, C.; Neumann, M. D.; Esser, N.; Fritze, S.; Witte, H.; Bläsing, J.; Dadgar, A.; Krost, A. Phys. Rev. B: Condens. Matter Mater. Phys. 2014, 90, 075203. (52) Tessarek, C.; Goldhahn, R.; Sarau, G.; Heilmann, M.; Christiansen, S. New J. Phys. 2015, 17, 083047. (53) Davydov, V. Y.; Averkiev, N. S.; Goncharuk, I. N.; Nelson, D. K.; Nikitina, I. P.; Polkovnikov, A. S.; Smirnov, A. N.; Jacobson, M. A.; Semchinova, O. K. J. Appl. Phys. 1997, 82, 5097−5102. (54) Schubert, E. F.; Goepfert, I. D.; Grieshaber, W.; Redwing, J. M. Appl. Phys. Lett. 1997, 71, 921−923. (55) Bussone, G.; Schäfer-Eberwein, H.; Dimakis, E.; Biermanns, A.; Carbone, D.; Tahraoui, A.; Geelhaar, L.; Haring, P. B.; Schülli, T. U.; Pietsch, U. Nano Lett. 2015, 15, 981−989. (56) Fauske, V. T.; Erlbeck, M. B.; Huh, J.; Kim, D. C.; Munshi, A. M.; Dheeraj, D. L.; Weman, H.; Fimland, B.-O.; van Helvoort, A. T. J. J. Microsc. 2016, 262, 183−188. (57) Sze, S. M.; Ng, K. K. Physics of Semiconductor Devices; John Wiley & Sons, Inc.: Hoboken, NJ, 2006. (58) Liu, Y.; Zhang, Z. Y.; Hu, Y. F.; Jin, C. H.; Peng, L. M. J. Nanosci. Nanotechnol. 2008, 8, 252−258. (59) Zhang, Z.; Yao, K.; Liu, Y.; Jin, C.; Liang, X.; Chen, Q.; Peng, L.M. Adv. Funct. Mater. 2007, 17, 2478−2489. (60) hyperspy: HyperSpy 0.8.1 http://hyperspy.org/. (61) Lumerical Solutions, Inc. http://www.lumerical.com/tcadproducts/fdtd/.

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Vertically Oriented Growth of GaN Nanorods on Si ... - ACS Publications

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