Letter pubs.acs.org/NanoLett
Epitaxy of GaN Nanowires on Graphene Vishnuvarthan Kumaresan,*,† Ludovic Largeau,† Ali Madouri,† Frank Glas,† Hezhi Zhang,‡ Fabrice Oehler,† Antonella Cavanna,† Andrey Babichev,‡,§ Laurent Travers,† Noelle Gogneau,† Maria Tchernycheva,‡ and Jean-Christophe Harmand† †
Laboratoire de Photonique et de Nanostructures (LPN), CNRS, Université Paris-Saclay, Route de Nozay, F-91460 Marcoussis, France ‡ Institut d’Electronique Fondamentale, UMR 8622 CNRS, University Paris Sud, University Paris-Saclay, 91405 Orsay cedex, France § ITMO University, St. Petersburg 197101, Russia S Supporting Information *
ABSTRACT: Epitaxial growth of GaN nanowires on graphene is demonstrated using molecular beam epitaxy without any catalyst or intermediate layer. Growth is highly selective with respect to silica on which the graphene flakes, grown by chemical vapor deposition, are transferred. The nanowires grow vertically along their c-axis and we observe a unique epitaxial relationship with the ⟨21̅1̅0⟩ directions of the wurtzite GaN lattice parallel to the directions of the carbon zigzag chains. Remarkably, the nanowire density and height decrease with increasing number of graphene layers underneath. We attribute this effect to strain and we propose a model for the nanowire density variation. The GaN nanowires are defect-free and they present good optical properties. This demonstrates that graphene layers transferred on amorphous carrier substrates is a promising alternative to bulk crystalline substrates for the epitaxial growth of high quality GaN nanostructures. KEYWORDS: Gallium nitride, GaN, nanowires, graphene substrate, TEM, photoluminescence
N
material with high thermal and electrical conductances. It can be synthesized at low cost on large areas.13 Furthermore, graphene films are transferable to almost any carrier substrate, including amorphous and/or flexible materials. This transfer can be performed before or after the deposition of an epitaxial structure on the graphene surface. Therefore, growth on graphene opens exciting perspectives for the fabrication of semiconductor devices without the need of a crystalline bulk wafer, where transparent contacts are also epitaxial seed layers and carrier substrate is transparent or flexible.14 Yet, published data on semiconductor epitaxy on graphene sheets is sparse. Detailed experimental studies have been reported for the specific cases of ZnO,15 GaAs,12 and InAs16 NWs. General considerations regarding various NWs of cubic or hexagonal phase have been reported to predict their possible orientations on the graphene lattice.12 Regarding GaN, there have been attempts to grow epilayers on graphite17 or graphene18 without or with intermediate material to promote GaN nucleation, such as ZnO nanowalls19 or AlN nucleation layer.20 The intermediate material is required as the GaN
itride semiconductor nanowires (NWs) are today widely explored as an active material for a large number of optoelectronic devices such as light-emitting diodes,1,2 photodetectors,3,4 or solar cells.5,6 In many cases, gallium nitride (GaN) NW growth is performed on bulk crystalline substrates, for example, silicon1,7,8 or sapphire.9,10 This results in vertical NWs of high quality in epitaxial relationship with their substrate. However, when devices are fabricated with such samples, the bulk substrates impose some of their characteristics that may not be adapted to the device functionality (e.g., low electrical or thermal conductivities, opacity, weight, rigidity, cost). Hence, there is a strong interest in developing alternative substrates without harm to the NW quality. A possible route is to use two-dimensional (2D) atomic layer materials transferred on a carrier substrate as seed layer for the NW growth. This effectively removes the epitaxial relationship between the carrier substrate and the NWs, which allows for a much wider range of material and a possible better fit between the host carrier characteristics and the device requirements. Growth on such a 2D material is often referred to as van der Waals heteroepitaxy.11 Although the exact bonding at the heterointerface between the crystal and the 2D material is still a matter of debate, it is not a usual covalent bonding. Recently, graphene sheets have been proposed as substrates to grow III−V semiconductor NWs.12 Graphene is a transparent and flexible © 2016 American Chemical Society
Received: April 7, 2016 Revised: July 8, 2016 Published: July 14, 2016 4895
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Figure 1. (a) Cross-section SEM image of GaN nanowires grown on a 1 cm2 patch of a graphene monolayer (ML). Inset is a 45°-titled view. (b) Top view SEM image showing four domains (marked with different background colors) with different NW in-plane orientations. Within each domain, all the NWs have the same sidewall facet orientation, as shown in the insets. (c) 2θ/ω scan showing diffraction peaks associated with GaN (0002) and Si (004) planes, indicating that GaN NWs grow along their c-axis. (d) Rocking curve measured for the GaN (0002) reflection. The FWHM (1.62°) characterizes the deviation of the NW orientation from verticality.
produce dense arrays of vertical NWs on pristine graphene sheets with a very good selectivity, that is, little to no NWs growing on the SiO2 surface. An epitaxial relationship between GaN NWs and graphene sheets is unambiguously established. We also observe a striking dependence of the nucleation rate upon the number of graphene layers. Results and Discussion. Graphene was grown by chemical vapor deposition (CVD) on copper substrates or nickel layers. Using different preparation conditions, we obtained either large polycrystalline flakes or small monocrystalline domains. Both types of graphene structures were transferred onto SiO2 (25 nm)/Si (100) carrier substrates using standard techniques (Supporting Information S1). Then, the graphene-on-silica substrates were loaded in the MBE chamber to grow GaN NWs. Neither buffer layer nor catalyst particles of any type were deposited before NW growth. We used a Ga flux corresponding to a planar growth rate of 0.62 monolayers of GaN per second and the N flow was adjusted to obtain an effective N/Ga flux ratio of 1.1. The plasma source was operated with a radio frequency power of 350 W. During growth, the background pressure was 8.3 × 10−6 Torr and the substrate temperature was maintained at 800 °C (see Supporting Information S2). Figure 1 shows scanning electron microscopy (SEM) views of GaN NWs grown on a large polycrystalline monolayer of graphene (about 1 cm2). All the NWs are vertical, as seen in the cross-section (Figure 1a). Their diameter is around 40 nm. Their in-plane orientation was analyzed by top view SEM observations. Typically, each isolated NW has six sidewall facets (m-planes) and more complex contours are observed when two or more NWs are merged. The facets reveal that NWs have the same in-plane orientation over large domains of several μm2, but that each large domain is randomly oriented in-plane. This is illustrated in Figure 1b where different zones of GaN NW inplane orientation can be observed. Each colored zone contains NWs with the same in-plane orientation and this orientation
nucleation rate on pristine and defect-free graphene is reported to be low.15,18−20 This is related to the sp2 orbital hybridization of graphene and its absence of dangling bonds. Consequently, GaN was observed to nucleate preferentially near ledges, kinks, or defects of the graphene surface.19 More specifically, growth of GaN NWs on graphitic surfaces was demonstrated by several groups. However, the NWs grow along several directions21,22 unless a buffer layer is used23 or unless graphene lies on a crystalline substrate,24 which yields no clear evidence of epitaxial relationship between the GaN and the graphitic surfaces. In particular, Heilmann et al. used single-layer graphene transferred onto sapphire substrates to obtain GaN micro- and nanostructures. In their case, highly vertical rods were obtained but these GaN structures grew predominantly on defective regions of graphene where the sapphire possibly helped the epitaxy. For comparison, this group transferred graphene on SiO2 and obtained in that case irregular structures with rather random growth directions. All these previous studies used metal organic chemical vapor deposition (MOCVD) for the growth of GaN. Two other studies have used plasma-assisted molecular beam epitaxy (PA-MBE). Araki et al.25 reported the formation of columnar GaN grains on single graphene layer transferred on Si(100) and Nakagawa et al. grew GaN NWs on thick graphite layers.26 From these two reports, PA-MBE seems more favorable than MOCVD to nucleate GaN NWs on graphitic surfaces because good quality GaN was obtained without intermediate layer or catalyst. Araki et al. found a majority of GaN grains having their c-axis normal to the substrate but their in-plane orientation was not reported. In Nakagawa’s work, the NWs orientation on graphite was not well-defined. In both cases, the exact epitaxial relationship between graphene and GaN was not discussed, calling for additional investigations. In the present work, we use the PA-MBE technique on monolayered and few-layered graphene transferred onto SiO2 as an alternative substrate for the growth of GaN NWs. We 4896
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Figure 2. (a) SEM top-view of isolated hexagonal flakes of graphene after growth on copper and transfer to a silica substrate. (b) SEM top-view of a single graphene flake with NWs grown selectively on the flake; inset showing the orientation of the hexagonal flake. (c) Top-view SEM image, zoomed at the edge of the flake showing the orientation of the NW facets; inset showing the orientation of the GaN NWs. (d) Schematics of the graphene honeycomb lattice and its zigzag edges at the periphery of the hexagonal flake and (e) schematics of the hexagonal cross-section of a NW with GaN lattice and {10 1̅ 0} sidewall facet.
Figure 3. (a) SEM bird’s eye view of the NWs grown near the central part of a pyramidal monocrystalline graphene flake. The flake consists of distinct terraces bounded by monolayer steps. From periphery to center, the graphene film thickness increases while the NW density and height both decrease. Insets show the NWs at the terrace boundaries of 2−3 (right) and 4−5 (left) graphene layers (b) NW density (squares), NW average height (dots), and width (triangles) versus graphene film thickness. The lines are guides for the eye.
rotates by a random angle when changing domains. We remark that the size of these domains is comparable to the graphene grain size. This strongly hints at a single epitaxial relationship between GaN and graphene. X-ray diffraction (XRD) was used to determine the NW growth direction on the same sample. A 2θ/ω analysis (Figure 1c) reveals the (0004) and (0002) diffraction peaks of GaN along with the (004) and (002) reflections from the (001) oriented Si substrate below the SiO2 carrier layer. The Si forbidden (002) reflection is related to multiple diffraction.27 Hence, as usually observed, the GaN NWs grow along their caxis. We quantified the verticality of the NW ensemble by recording an XRD rocking curve for the GaN (0002) reflection (Figure 1d). A macroscopic 1 × 5 mm2 area was irradiated by the X-rays for this measurement. The (0002) reflection has a full width at half-maximum (FWHM) of 1.62°. Though this value demonstrates a very good verticality of the NW ensemble, it indicates a residual misalignment. This can be attributed to an imperfect transfer process of the graphene film that could
present cracks, undulations, or folds affecting the verticality of some NWs. Nevertheless, the average deviation is below 2°, which is a major improvement from the previous literature data.21−23,26 It is important to note that the presence of the 25 nm thick SiO2 intermediate carrier layer guarantees that the observed NW alignment is not induced by the crystalline Si substrate. This SiO2 thickness does not appear critical for GaN NW alignment since vertical NWs were obtained also on 1 ML graphene transferred on 50 and 300 nm SiO2 layers (see Supporting Information S3). To investigate further the in-plane alignment of NWs grown on graphene and to determine the exact epitaxial relationship between these two materials, we grew GaN NWs on isolated monocrystalline graphene flakes. The graphene samples were fabricated on copper foils by pulsed-CVD using alternate cycles of CH4 (20 sccm, 5 s) and H2 (960 sccm, 55 s), a substrate temperature of 1000 °C, and a reactor pressure of 25 mbar (see Supporting Information S1). Using these conditions, graphene forms as separate hexagonal domains that extend laterally. The 4897
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Figure 4. (a) SEM images of the TEM sample preparation: (left) top view image of the center of the graphene flake after carbon deposition; (middle) 52° tilted view after Pt deposition; (right) after bulk-out, lift-out, and sticking to the Cu TEM grid. (b) GaN NWs cross sectional TEM images taken across 5, 4, 3, and 2 layers of graphene and (c) corresponding selected area diffraction patterns showing a single in-plane orientation of the NWs.
few-layered graphene domains,28 we can reasonably assume that the flake is pyramidal with monolayer steps, that is, the number of graphene monolayers decreases by one unit every step from center to periphery. Each step forms a hexagonal contour that is perfectly reflected by abrupt changes of NW density. Close to the periphery, the flake consists of a single graphene layer with a very high NW density (340 NWs/μm2). In contrast, the NW density is much lower (190 NWs/μm2) in the 5 ML central terrace. The dependences of NW height and density are summarized in Figure 3b and will be discussed later. To unambiguously determine the relative crystal orientation between GaN NWs and graphene at each staggered terrace, we have prepared a transmission electron microscopy (TEM) sample from a pyramidal graphene flake with six terraces. To this aim, we prepared a thin foil for transmission electron microscopy (TEM); the sample was coated with benzocyclobutene (BCB) and carbon before depositing a Pt strip perpendicularly to a series of terrace edges of the graphene domain. The strip covers regions with 2, 3, 4, 5, and 6 graphene layers. A thin lamella below the strip was extracted and mounted for cross sectional observation (Figure 4a). The different NW heights along the lamella appear clearly on the TEM image and correspond to the number of graphene layers indicated in Figure 4b. Selective area electron diffraction (SAED) patterns are shown for NW ensembles from each zone (Figure 4c). All the patterns are similar. They reveal the wurtzite structure of the NWs and demonstrate that these all have the same in-plane orientation, irrespective of the number of graphene layers underneath. This indicates that all the graphene terraces should also have a common orientation, as strongly suggested by their concentric and parallel edges (Figure 4a). The in-plane orientation of the first graphene layer is thus preserved when the pyramidal stack forms. Some SAED patterns present curved spots indicative of some deviation in the vertical alignment of the NWs. This might result from the
process is stopped well before the coalescence of these domains and regular hexagonal monocrystalline graphene flakes are obtained. Figure 2a shows a SEM top view of such graphene flakes. When a flake is sufficiently large, it can consist of several concentric areas with a number of monolayers decreasing from center to periphery, as will be seen later in Figure 3. The morphology of such multilayered graphene flakes is often reported and has been studied in detail for similar growth conditions.28 It is important to note that the edge of the graphene flakes grown on copper foils is well established to be of zigzag type.29,30 After transferring the isolated graphene flakes onto a SiO2/Si carrier substrate, GaN NW growth was performed under the same conditions as that of the previous sample. The image in Figure 2b shows the top view SEM of the sample after growth. The growth is highly selective: NWs are present on the graphene flake but not on the SiO2 carrier substrate. We determined the orientation of the NW facets relative to the edge of the graphene flakes (Figure 2c). As expected from our previous observations, all of the NWs grown on a single flake have the same in-plane orientation, their lateral facets corresponding to the {1010̅ } planes. Schematics of the GaN lattice orientation is shown in Figure 2e. The hexagonal edge of the graphene flake (zigzag type) gives the crystalline orientation of the graphene layer (Figure 2d). From this, it is straightforward to deduce the relative in-plane orientation of the two materials: the ⟨211̅ 0̅ ⟩ directions of the wurtzite GaN lattice are parallel to the directions of the zigzag chains of the honeycomb carbon lattice. This in-plane relationship is confirmed by the study of a larger monocrystalline flake with up to five staggered monolayers at the center of the flake. Figure 3a shows the central region of such a pyramidal flake after NW growth. Strikingly, the NW height and density both decrease in a step-like fashion from the center of the flake toward its periphery. Referring to detailed studies of similar 4898
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Figure 5. Schematics of possible epitaxial relationships between graphene and the first atomic plane of GaN NWs. The black and red diamonds are primitive in-plane cells of graphene and GaN, respectively. The blue diamond is the cell of coincidence of the two lattices strained to accommodate a −3.1% misfit. The two possible arrangements are obtained by positioning some atoms of the GaN plane in H and T sites (a) or B sites (b) of the graphene lattice. These three sites (top, bridge, hollow) are recalled in (c).
nonideal transfer process of the graphene domain or from strain related to the fabrication of the lamella. Indexation of the diffraction spots indicates that the GaN [11̅00] direction is parallel to the long axis of the lamella, that is, perpendicular to the zigzag edges of the graphene terraces. This is consistent with the in-plane orientation that we first deduced from the observation of NW facets by SEM (Figure 2). Therefore, the entire NW ensemble grew coherently on the whole flake with the in-plane relationship previously indicated. The relative orientation of the two lattices does not correspond to that anticipated by Munshi et al.12 Let us speculate about the possible epitaxial relationship. The in-plane lattice parameters are 0.24682 nm for graphene and 0.3189 nm for GaN. The misfit is much too large (29.2%) to be fully accommodated by coherent strain. On the other hand, one can notice a good coincidence between the two lattices every 4 unit cells of graphene and 3 unit cells of GaN. Indeed, the residual misfit associated with this coincidence is of only −3.1%, which seems compatible with strain accommodation. If this is the case, the first atomic plane of GaN (which could be either a Ga or N plane) could be positioned on the graphene lattice as depicted in Figure 5. One possible configuration consists of a supercell containing one III or V atom in H site, two in T site, and six at positions close to B sites (Figure 5a). Another possibility (less likely) is obtained by placing one atom on a B site with the other eight being at intermediate positions (Figure 5b). We now examine the crystalline structure at the NW base. The TEM observations of Figure 6 show a perfect wurtzite structure, free of defect. This TEM analysis could not reveal the atomic arrangement of the graphene layers below the NWs, probably because of surface roughness of the SiO2 layer underneath. Let us now discuss what the possible origin of the variation of NW density with the number of staggered graphene layers. It is interesting to compare our observations with the reports of Au deposition on graphene, which show the formation of Au nanostructures with a density decreasing with increasing number of graphene layers.31,32 For these Au nanostructures, the decrease of density is accompanied by an increase of the nanostructure volume so that the total deposited Au volume remains constant independently of the terrace. The results were interpreted in terms of adatom diffusivity decreasing strongly
Figure 6. {112̅0} zone axis STEM bright field images. (a) The NW base and sidewalls (b) zoom on the NW−graphene−silica interface showing the base of the nanowires, showing a perfect wurtzite structure.
for the smaller number of graphene layers. This interpretation cannot apply to our case since the total volume of GaN per unit area varies strongly with the number of graphene layers (see Figure 3b). Another explanation could result from the Ehrlich− Schwoebel barrier associated with the step edge of each terrace. This would create a net flux of adatoms from the thick to the thin parts of the flake, inducing an easier and quicker nucleation on the latter. To test this hypothesis, we investigated the impact of the number of graphene layers using large patches of uniform thickness. We transferred two continuous polycrystalline patches of about 1 cm2 each onto the same SiO2/Si carrier substrate. The first patch is a single graphene layer, the second consists of few-layered graphene with a nominal average thickness equal to four monolayers. We used our standard PA-MBE conditions to elaborate NWs in a single growth run on this sample. Similarly to the small pyramidal flakes, the NWs are much denser and taller on the monolayer patch (Figure 7). This rules out a possible effect related to Ehrlich−Schwoebel barriers because we observe the same density differences for the much larger terraces of this latter sample. 4899
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Figure 7. SEM tilted views of GaN NWs grown on (a) the monolayer patch and (b) few-layered graphene patch from the same growth run. Insets show corresponding top-view SEM images.
Figure 8. NW density as a function of number of underlying monolayers of graphene. The curve gives the best fit of the experimental data (dots, from Figure 3b) using eq 2. A continuous curve is given as a guide, although the density is only defined for integer values of n.
We then propose and test a third hypothesis, based on elasticity. In accordance with classical nucleation theory, we shall assume that the NWs nucleate from a sea of adatoms; the adatom surface density increases with time until some critical value is reached and enables the formation of one monolayer high nuclei. Considering the epitaxial relationship that we propose, the formation of a coherent GaN nucleus on graphene must induce some strain (possibly 3.1%). A 2D GaN nucleus of 1 monolayer has a thickness of the same order as that of the flake. Then, contrary to the case of thick substrates, the epitaxial strain partitions between graphene and nucleus, and this partitioning depends on the graphene thickness. Here, our aim is only to assess the magnitude of this effect and its possible influence on nucleation and hence on NW density. To this end, we adopt a very simple model. We assume flat 2D nuclei and a compliant graphene flake (no strain is induced by the carrier substrate) and neglect strain relaxation at the nucleus edges. In regard to elasticity, we thus treat the nucleus as laterally infinite. We also take both materials as elastically isotropic with the same Young’s modulus E and Poisson ratio ν. The elastic energy stored in the system per unit volume of nucleus is then easily calculated as
we =
n E 2 ϵ0 n+ 11− ν
mention the extracted value of the ratio w∞/Δμ of these quantities, namely 1.18, which means that the elastic energy per pair is of the same order as the difference of chemical potential between adatoms and solid. Assuming a misfit of 3.1%, we estimate that Δμ ≃ 90 meV/GaN pair. This value seems entirely plausible for our PA-MBE conditions. At low nucleation probability, the nucleation events are delayed and spread in time. This probably explains why the NWs are shorter. Finally, the optical properties of NW arrays on monolayer graphene were probed by photoluminescence (PL) spectroscopy. The PL characterization was performed on NW ensembles using a frequency-doubled continuous-wave Ar2+ ion laser (λ = 244 nm). The laser was focused on the NW array into a spot of about 5 μm by means of a 20× UV microscope objective with a 0.4 numerical aperture. The PL was analyzed in a 460 mm focal length grating spectrometer and a CCD camera, with a spectral resolution roughly equal to 1 nm. Figure 9 shows the spectrum collected from a NW array grown on a 1 cm2 patch of 1 ML thick graphene. PL for GaN NWs on graphene domains was also measured and the spectral response was the same (see Supporting Information S4). The spectrum presents the GaN near band edge (NBE) emission peaked at 3.472 eV with a FWHM of 14 meV. Gaussian fit of this peak allows to identify three spectral contributions peaked at 3.472, 3.45 and 3.42 eV with FWHMs of 14, 24, and 31 meV respectively. The 3.472 eV peak has the highest intensity and is attributed to neutral donor-bound exciton transition (D0XA). It is relatively wide but comparable to that we obtained for GaN NWs grown on conventional Si(111) substrates.34 Peak broadening could be due to a large residual doping in our MBE growth chamber. The 3.45 eV sub-bandgap contribution is often reported in GaN NWs and attributed mainly to surface defects.35−37 Following the recent in-depth analyses of this line,38 it can be attributed to inversion domain boundaries, which have been reported to appear in GaN NWs grown on Si substrates by MBE. This peak has a low intensity in the studied NWs on graphene, which indicates that these defects remain scarce. In fact, this peak is weaker than for the NWs grown on Si(111),34 a possible sign that NWs grown on graphene contain less inversion domain boundaries than those grown on Si(111). Finally, the 3.42 eV weak sub-bandgap emission has been attributed in previous reports to excitons bound to planar defects such as stacking faults.39,40 Further investigations will be carried out to verify this attribution. A similar peak is also observed in our GaN NWs grown on Si(111) substrates.34
(1)
with ϵ0 as the in-plane misfit between graphene and GaN and n as the number of graphene layers. Using eq 1, the nucleation probability then writes ⎡ −αγ 2 ⎢ P = P0 exp⎢ n ⎢⎣ kBT Δμ − n + 1 w∞
(
)
⎤ ⎥ ⎥ ⎥⎦
(2)
where kB is the Boltzmann constant, T is the temperature, α is a geometric constant, γ is the edge energy of the GaN nucleus, Δμ is the difference of chemical potential of a (Ga,N) pair between adatom state and solid state in the absence of elastic effects, and P0 is a quantity that varies only slowly with Δμ compared to the exponential term.33 Elastic effects oppose nucleation by diminishing Δμ of the elastic energy per Ga−N pair (obtained by scaling eq 1) with w∞ as the limit of this energy for an infinite number of graphene layers (n → ∞). The last step is to assume that the NW density is proportional to the nucleation probability before formation of the first NWs and that the adatom supersaturation (Δμ) is the same on all terraces. Figure 8 shows that eq 2 fits the measured ndependence of the NW density very well. From this fit, we extract the absolute values of kBTΔμ/(αγ2) and of w∞/(αγ2) . This will be discussed in detail elsewhere. Here, we only 4900
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Figure 9. (a) Photoluminescence spectrum of GaN NWs on a 1 cm2 patch of 1 ML graphene substrate at 4 K. (b) Gaussian fitting of the near-bandedge emission evidencing three spectral contributions at 3.42, 3.45, and 3.472 eV.
domain monolayer and multilayer graphene substrates. H.Z performed the PL measurements. V.K., L.L., and J.-C. H. proposed the relationship of coincidence between graphene and GaN. F.G. formulated the model for the NW density scaling with respect to the number of monolayers of graphene. All other authors contributed equally in analyzing and writing the results.
Therefore, the optical quality of GaN NWs grown on grapheneon-amorphous templates is not degraded with respect to the GaN NWs on a crystalline Si substrate. Conclusion. In conclusion, we have shown that CVDgrown graphene is an excellent substrate for GaN NW growth by PA-MBE. Epitaxy is demonstrated on large or small graphene flakes transferred onto amorphous SiOx carrier layers. In our conditions, growth is highly selective; NWs nucleate on graphene but not on the surrounding silica. The nanowires grow vertically along their c-axis and we observe a unique epitaxial relationship with the ⟨21̅1̅0⟩ directions of the wurtzite GaN lattice parallel to the directions of the carbon zigzag chains. We propose that the two materials are strained to accommodate a −3.1% in-plane misfit ensuring a certain lattice coincidence. Strikingly, the number of graphene layers has a strong impact on GaN nucleation; the NW density and height decrease rapidly when this number increases from 1 to 5. This original observation is interpreted as a result of the interfacial strain that at NW nucleation partitions between graphene and GaN nucleus. The higher number of graphene layers, the larger elastic energy stored in the system, which elevates the nucleation barrier. Finally, the PL of the GaN NWs grown on graphene shows a near band edge emission dominated by the D0XA transition. The optical quality is comparable to that of GaN NWs grown on crystalline Si substrate. This successful epitaxial growth of GaN NWs on graphene seems particularly promising for the development of flexible devices.
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Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work has received financial support by the French “Agence Nationale de la Recherche” (ANR) through “GaNeX” (ANR11- LABX-2014) and “NanoSaclay” (ANR-10-LABX-0035) programs. We also acknowledge ANR “Investissement d’Avenir” program (TEMPOS project no. ANR-10-EQPX-50) for having funded the acquisition of the NANOTEM platform (Dual beam FIB-FEG FEI SCIOS system and TEM-STEM FEI Titan Themis) used in this work. A.B. acknowledges the support of RFBR (project numbers 14-02-31485, 15-02-08282) and of Scholarship of the President of the Russian Federation (Grant SP-4716.2015.1).
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ABBREVIATIONS GaN, gallium nitride; NW, nanowire; SEM, scanning electron microscopy; XRD, X-ray diffraction; FWHM, full width at halfmaximum; TEM, transmission electron microscopy; SAED, selective area electron diffraction; PL, photoluminescence; 2D, two dimensional; PA-MBE, plasma-assisted molecular beam epitaxy; CVD, chemical vapor deposition; MOCVD, metalorganic chemical vapor deposition; BCB, benzocyclobutene
ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.nanolett.6b01453. Synthesis and transfer of monolayered and multilayered graphene domains (S1); GaN nanowire growth by plasma-assisted molecular beam epitaxy (S2); thickness of SiO2 carrier layer (S3); photoluminescence for GaN NWs on graphene domains (S4).(PDF)
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REFERENCES
(1) Guo, W.; Zhang, M.; Banerjee, A.; Bhattacharya, P. Nano Lett. 2010, 10, 3355. (2) Hersee, S. D.; Fairchild, M.; Rishinaramangalam, A. K.; Ferdous, M. S.; Zhang, L.; Varangis, P. M.; Swartzentruber, B. S.; Talin, A. A. Electron. Lett. 2009, 45, 75. (3) Rigutti, L.; Tchernycheva, M.; De Luna Bugallo, A.; Jacopin, G.; Julien, F. H.; Zagonel, L. F.; March, K.; Stephan, O.; Kociak, M.; Songmuang, R. Nano Lett. 2010, 10, 2939. (4) González-Posada, F.; Songmuang, R.; Den Hertog, M.; Monroy, E. Nano Lett. 2012, 12, 172. (5) Tian, B.; Kempa, T. J.; Lieber, C. M. Chem. Soc. Rev. 2009, 38, 16. (6) Tang, Y. B.; Chen, Z. H.; Song, H. S.; Lee, C. S.; Cong, H. T.; Cheng, H. M.; Zhang, W. J.; Bello, I.; Lee, S. T. Nano Lett. 2008, 8, 4191. (7) Calarco, R.; Meijers, R. J.; Debnath, R. K.; Stoica, T.; Sutter, E.; Lüth, H. Nano Lett. 2007, 7, 2248.
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Author Contributions
V.K., M.T., and J.-C.H. conceived the experimental work. V.K. carried out the nanowire growth experiment and SEM characterization. L.L. performed the TEM and XRD characterisations. A.M. and A.C. grew and transferred the monocrystalline graphene domains. A.B. and M.T. supplied the multi4901
DOI: 10.1021/acs.nanolett.6b01453 Nano Lett. 2016, 16, 4895−4902
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Nano Letters (8) Largeau, L.; Galopin, E.; Gogneau, N.; Travers, L.; Glas, F.; Harmand, J.-C. Cryst. Growth Des. 2012, 12, 2724. (9) Hersee, S. D.; Sun, X.; Wang, X. Nano Lett. 2006, 6, 1808−1811. (10) Wang, G. T.; Talin, A. A.; Werder, D. J.; Creighton, J. R.; Lai, E.; Anderson, R. J.; Arslan, I. Nanotechnology 2006, 17, 5773. (11) Hong, Y. J.; Lee, C.-H. Semiconductor Nanowires I - Growth and Theory; Elsevier: New York, 2015. (12) 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. (13) Kobayashi, T.; Bando, M.; Kimura, N.; Shimizu, K.; Kadono, K.; Umezu, N.; Miyahara, K.; Hayazaki, S.; Nagai, S.; Mizuguchi, Y.; Murakami, Y.; Hobara, D. Appl. Phys. Lett. 2013, 102, 023112. (14) Chung, K.; Lee, C.-H.; Yi, G.-C. Science 2010, 330, 655. (15) Kim, Y.-J.; Lee, J.-H.; Yi, G.-C. Appl. Phys. Lett. 2009, 95, 213101. (16) Hong, Y. J.; Yang, J. W.; Lee, W. H.; Ruoff, R. S.; Kim, K. S.; Fukui, T. Adv. Mater. 2013, 25, 6847. (17) Miura, A.; Shimada, S. Mater. Res. Bull. 2006, 41, 1775. (18) 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. (19) Baek, H.; Lee, C. H.; Chung, K.; Yi, G. C. Nano Lett. 2013, 13, 2782. (20) Gupta, P.; Rahman, A. A.; Hatui, N.; Gokhale, M. R.; Deshmukh, M. M.; Bhattacharya, A. J. Cryst. Growth 2013, 372, 105. (21) Yuan, F.; Liu, B.; Wang, Z.; Yang, B.; Yin, Y.; Dierre, B.; Sekiguchi, T.; Zhang, G.; Jiang, X. ACS Appl. Mater. Interfaces 2013, 5, 12066. (22) Park, J. B.; Kim, N.-J.; Kim, Y.-J.; Lee, S.-H.; Yi, G.-C. Curr. Appl. Phys. 2014, 14, 1437. (23) Chung, K.; Beak, H.; Tchoe, Y.; Oh, H.; Yoo, H.; Kim, M.; Yi, G.-C. APL Mater. 2014, 2, 092512. (24) Heilmann, M.; Sarau, G.; Göbelt, M.; Latzel, M.; Sadhujan, S.; Tessarek, C.; Christiansen, S. Cryst. Growth Des. 2015, 15, 2079. (25) 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. (26) Nakagawa, S.; Tabata, T.; Honda, Y.; Yamaguchi, M.; Amano, H. Jpn. J. Appl. Phys. 2013, 52, 08JE07. (27) Hwang, B.-H. J. Phys. D: Appl. Phys. 2001, 34, 2469. (28) Robertson, A. W.; Warner, J. H. Nano Lett. 2011, 11, 1182. (29) Yu, Q.; Jauregui, L. A.; Wu, W.; Colby, R.; Tian, J.; Su, Z.; Cao, H.; Liu, Z.; Pandey, D.; Wei, D.; Chung, T. F.; Peng, P.; Guisinger, N. P.; Stach, E. A.; Bao, J.; Pei, S.-S.; Chen, Y. P. Nat. Mater. 2011, 10, 443. (30) Luo, Z.; Kim, S.; Kawamoto, N.; Rappe, A. M.; Johnson, A. T. C. ACS Nano 2011, 5, 9154. (31) Liu, L.; Chen, Z.; Wang, L.; Polyakova Stolyarova, E.; Taniguchi, T.; Watanabe, K.; Hone, J.; Flynn, G. W.; Brus, L. E. J. Phys. Chem. B 2013, 117, 4305. (32) Zhou, H.; Qiu, C.; Liu, Z.; Yang, H.; Hu, L.; Liu, J.; Yang, H.; Gu, C.; Sun, L. J. Am. Chem. Soc. 2010, 132, 944. (33) Glas, F.; Harmand, J.-C.; Patriarche, G. Phys. Rev. Lett. 2010, 104, 135501. (34) Kumaresan, V.; Largeau, L.; Oehler, F.; Zhang, H.; Mauguin, O.; Glas, F.; Gogneau, N.; Tchernycheva, M.; Harmand, J.-C. Nanotechnology 2016, 27, 135602. (35) Brandt, O.; Pfüller, C.; Chèze, C.; Geelhaar, L.; Riechert, H. Phys. Rev. B: Condens. Matter Mater. Phys. 2010, 81, 045302. (36) Sam-Giao, D.; Mata, R.; Tourbot, G.; Renard, J.; Wysmolek, a.; Daudin, B.; Gayral, B. J. Appl. Phys. 2013, 113, 043102. (37) Lefebvre, P.; Fernández-Garrido, S.; Grandal, J.; Ristić, J.; Sánchez-García, M.-A.; Calleja, E. Appl. Phys. Lett. 2011, 98, 083104. (38) Auzelle, T.; Haas, B.; Den Hertog, M.; Rouvière, J.-L.; Daudin, B.; Gayral, B. Appl. Phys. Lett. 2015, 107, 051904. (39) Forsberg, M.; Serban, A.; Poenaru, I.; Hsiao, C.-L.; Junaid, M.; Birch, J.; Pozina, G. Nanotechnology 2015, 26, 355203.
(40) Lähnemann, J.; Jahn, U.; Brandt, O.; Flissikowski, T.; Dogan, P.; Grahn, H. T. J. Phys. D: Appl. Phys. 2014, 47, 423001.
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DOI: 10.1021/acs.nanolett.6b01453 Nano Lett. 2016, 16, 4895−4902
