GaN nanowires with periodic serrated morphology have been synthesized on Si substrate by Au-catalyzed vapor-liquid-solid growth mode. The presence of Mn vapor during growth process has been found to enhance the production and quality of serrated GaN nanowires, without introducing dopants. We have performed photoluminescence and Raman spectral measurements on nanowires with different levels of serration. Temperature dependent photoluminescence revealed a broad yellow-green and red luminescence in the samples. Room temperature Raman spectra exhibits disorder-activated phonon mode at ~670 cm−1, in addition to E2(high) and A1(LO) modes of GaN. Further investigation of Raman spectra revealed the presence of tensile stress in the GaN nanowires when Mn vapor is present during the growth process. The dependence of the optical properties on the morphology of GaN nanowires shows that they can be tuned by initial synthesis conditions.
© 2014 Optical Society of America
Wurtzite GaN is a wide band gap material with direct band gap of 3.4 eV at 300 K. Robustness against radiation damage, high melting temperature, physical hardness and excellent thermal conductivity are the salient features of GaN which makes it a favorable candidate for short wavelength light emitting devices and high power devices [1,2]. Moreover, GaN can be easily doped with Si or Mg for n or p type conductivity respectively and hence the carrier concentration can be easily tuned at will [3,4]. Alloying GaN with In or Al also allows band gap engineering which most importantly spans the entire solar spectrum (0.7-3.4 eV) [5,6]. Hence, over the years, GaN has emerged as one of the most desirable material for optoelectronics. Often GaN is doped with magnetic impurity, such as Mn to realize a dilute magnetic semiconductor. While the Curie temperature of Mn doped InAs and Mn doped GaAs is ~25 K and ~110 K, respectively [7,8], it was theoretically predicted that the same for Mn doped GaN can exceed room temperature [9,10]. Many research groups have experimentally validated the theoretical predictions [11–13]. In fact, a high doping of Mn in GaN is expected to exhibit room temperature ferromagnetic behavior and might be a desirable candidate for spintronic applications. Experimental realization of group III-nitride nanowire photovoltaic devices , demonstration of high electron mobility transistors in GaN/AlN/AlGaN nanowire heterostructure  and field effect transistors based on individual GaN nanowire  have shown great potential of GaN nanowire-based devices for future optoelectronic and high temperature microelectronic devices and have further surged the research interest in GaN nanowire.
In the present work, we report the optical properties of recently realized periodically serrated GaN nanowire . These serrated GaN nanowires were synthesized with the introduction of MnO2 powder in the reaction source along with Ga2O3 in horizontal quartz tube inside CVD furnace under optimized conditions. It is interesting to note that the Au-catalyzed vapor-liquid-solid growth of GaN nanowires by using MnO2 powder does not incorporate any Mn doping in GaN but does introduce some defect states which in turn gives rise to different optical and electrical signatures. The serrated GaN nanowires, owing to large surface to volume ratio, are expected to show better sensitivity in sensor applications.
GaN thin films are routinely grown by molecular beam epitaxy (MBE). For single crystal GaN nanowires, chemical vapor deposition and metal-organic vapor phase epitaxy are frequently used. GaN nanowires in the present work have been grown by chemical vapor deposition process which is governed by vapor-liquid-solid mechanism. Mostly, a catalytic liquid alloy phase is introduced so as to speed up reaction rates. In view of the same, a thin film of Au of thickness ~10 nm, deposited by e-beam evaporation, has been used as a catalyst on Si(100) substrate. Inside a horizontal quartz tube, from upstream to downstream, MnO2 powder (source of Mn), Ga2O3 powder (source of Ga) and Au coated Si substrate were placed and a pressure of 100 torr is maintained. The quartz tube was heated to the temperature of ~960°C, a typical growth temperature for GaN nanowires. A gas of ammonia (as reaction source of N) and a gas of hydrogen (to decompose Ga2O3) were then introduced in quartz tube at a flow rate of 50 sccm and 30 sccm, respectively. The mass of Ga2O3 was 0.5g and two samples with different mass of MnO2, viz., 0.2 g and 0.4 g were synthesized. During preheating and cooling processes, Ar was flowing in furnace. Typical growth time lasts for about 1 hour. In this work, three samples are investigated. Sample 1 is GaN nanowires with no MnO2 (with 0.5 g of Ga2O3 as source), sample 2 is GaN nanowires grown with smaller quantity of Mn (0.5 g of Ga2O3 and 0.2 g of MnO2 as sources) and sample 3 is GaN nanowires with larger quantity of MnO2 (0.5 g of Ga2O3 and 0.4 g of MnO2).
Morphology of the nanowires in the three samples is imaged by scanning electron microscopy (SEM). SEM images in Fig. 1 show that the increasing presence of MnO2 in the growth chamber results in increased density of serrated nanowires. Structural characterization was performed by X-ray diffraction (XRD). Temperature dependent photoluminescence (PL) measurements were performed by exciting with 325 nm line of a He-Cd laser. Micro-Raman spectroscopy measurements were performed at room temperature using 632 nm laser line and a 10x objective with about 1 µm spatial and 2 cm−1 spectral resolution. Cr/Au contacts were deposited by e-beam lithography on individual GaN nanowires (both straight and serrated) which were previously released from solution onto a Si/SiO2 substrate (see  for more details). Two-probe electrical measurements were subsequently carried out inside a probe station (Janis ST-500).
3. Results and discussion
Representative SEM images of the three samples are shown in Fig. 1. As seen from the SEM of sample-1 (Fig. 1(a)), GaN nanowires are densely formed and are uniformly dispersed over the Si substrate. Figure 1(b) clearly shows the existence of serrated and straight (smooth sidewall) morphology for as-grown nanowires (sample-1). Figures 1(c) and 1(d) display the SEM images of sample-2 and sample-3, respectively. It can be inferred from the SEM images that introduction of MnO2 tend to alter the morphology of GaN nanowire. Although all the three samples show the coexistence of serrated and straight GaN nanowires, it is evident that the higher the content of MnO2 source during growth process higher is the density of serrated nanowires. Recently, by adopting a similar growth procedure for the GaN nanowire (without MnO2 source) in CVD , it was discovered that the growth parameter window for serrated morphology can be enhanced by simultaneously meeting two criteria: a) large Au diameter and b) large Ga/N ratio. It was inferred that the kinetic frustration at the growth front induced by the energetic and geometric constraints leads to the realization of diameter modulated nanowire along the length. The growth mechanism of impurity (MnO2) introduced morphological change in GaN nanowire, however, needs further investigation and a possible explanation will be presented below after the structural and optical properties.
The X-ray diffraction spectra of sample-1, sample-2 and sample-3 are displayed in Figs. 2(a)-2(c), respectively. All the peaks are indexed to the hexagonal wurtzite structure of GaN. Reflection corresponding to Au (111) and Au (200) peaks, arising from the tip of nanowires were also observed in the XRD spectra. However, no trace of Mn could be detected which was confirmed by EDAX measurements as well. The signal to noise ratio in XRD measurements seems better when MnO2 is present in the growth chamber.
Various point defect-related radiative transitions in GaN samples are believed to give rise to luminescence bands. The variation of PL emission of GaN sample with temperature is plotted in Fig. 3(a).Sample-1 showed PL emission peaks at ~440 nm and ~500 nm, with the emission at longer wavelength side stronger than the shorter wavelength. The room temperature PL spectrum has considerably low intensity than that at low temperature. The emission at ~440 nm is frequently referred to as the blue emission and is attributed to the transition from either conduction band or a shallow donor to a relatively deep acceptor, whose ionization energy is about 0.34-0.4 eV . As seen from Fig. 3(a), the intensity of the blue emission increases with temperature, however, at room temperature the fine structure at ~440 nm almost vanishes. Similar observation has been reported earlier [19,20] where the increase in PL with temperature was argued to be due to redistribution of holes, resulting from exciton dissociation, between the acceptor levels. At room temperature, on the other hand, the holes in acceptor level thermalize to the valence band giving rise to a lower emission at ~440 nm [19,20]. The blue emission becomes negligible in case of GaN samples grown in presence of MnO2, as is evident from Fig. 3(b) and Fig. 3(c) which display the PL spectra of sample-2 and sample-3, respectively.
The emission at ~500 nm and the asymmetry towards longer wavelength side in sample-1 may be the outcome of the combination of multiple emission bands viz., green emission (~500 nm), yellow emission (~560 nm) and red emission (~590 nm). The origin of broad yellow-green band in GaN has been a debatable subject and consequently is an area of intensive research [19,21,22]. Nevertheless, it is widely accepted that these bands are the signature of the presence of multiple charge state acceptors . These multiple charge state acceptors are complex involving a Ga vacancy (VGa) and a shallow oxygen donor on a N site (ON). The green luminescence (~500 nm) and the yellow luminescence (~550 nm) originate from the two charge states of the same defect and are assigned to -/0 and −2/- states of the VGaON complex, respectively. Reports on red luminescence, however, are scarcely studied and therefore less common in undoped GaN as compared to yellow and green luminescence. There is a considerable scatter in the reported peak position of red luminescence in GaN (1.5 eV-2 eV) . It is also observed that the red luminescence is independent of temperature and it originates from transition from a shallow donor (at low temperature) or the conduction band (at elevated temperature) to a deep acceptor level [19,24]. The prominent emissions at higher energy seen at lower temperatures are less pronounced at room temperature in the PL of sample-1. The emission at room temperature is broad and beyond 600 nm there is no appreciable change in the spectra in the 4 K to 300 K temperature range. The PL spectra of sample-2 and sample-3 are broad and symmetric and have negligible blue emission, as is evident from Fig. 3(b) and Fig. 3(c), respectively. A broad PL emission centered at ~590 nmis evident, in addition to a hump at ~500 nm in both sample-2 and sample-3. The intensity does not differ much with change in temperature for sample-3 (Fig. 3(c)). However, a slight decrease in PL emission is registered for sample-2. To compare, low temperature (4 K) PL spectra of all the three samples that are normalized with respect to the maximum PL signal are plotted in Fig. 3(d). The low temperature spectra for the three samples show a significant difference. Sample-2 and sample-3, exhibit similar behavior. However, sample-1 shows pronounced emission at 440 nm and 500 nm at 4 K. One of the possible reasons behind the existence of luminescence at high energy in sample-1 may be attributed to the difference in the degree of confinement in the nanowires. It may be noted that the diameter of the regular wire (abundant in sample-1) is ~75 ± 25 nm and the outer diameter of the serrated nanowires (abundant in sample-3) is 250 ± 150 nm. Consequently, the PL at higher energies is visible in Sample 1 but not in the other two samples. From Fig. 3(e), it is clear that all the three samples exhibit defect related emission at room temperature and the broad spectrum of sample-3 indicates that the defect density increases with increase in MnO2 impurity level. The band edge luminescence is expected at about 360 nm in GaN, however, we do not see this in our samples. This could be due to fast carrier trapping at defect sites but requires further studies.
Raman scattering is an inelastic scattering phenomenon which can provide a deeper insight about molecular vibrations of the sample under investigation. Near zone centre, single crystal GaN possesses 8 phonon modes (A1 + 2B1 + E1 + 2E2). Out of these, the A1, E1 and 2E2 modes are Raman active and the two B1 modes are Raman inactive. The A1 and E1 modes have longitudinal optical (LO) and transverse optical (TO) splitting. The Raman spectra of GaN nanowires with and without MnO2 content during growth are displayed in Fig. 4.The peaks at ~568 cm−1 and ~730 cm−1 are indexed to the E2(high) and A1(LO) modes of GaN, respectively . The higher intensity of E2(high) mode reveals hexagonal wurtzite structure of GaN nanowires. The band at 420 cm−1 is the zone boundary phonon mode . A sharp peak at ~521 cm−1 observed in GaN sample is due to Si substrate. An additional small signal at ~670 cm−1 is seen in the case of GaN samples with MnO2 content (sample-2, Fig. 4(b) and sample-3, Fig. 4(c). Inset to Fig. 4). The low intense 670 cm−1 band in GaN has been reported to be the defect-induced phonon mode . It can thus be inferred that MnO2 incorporation might have induced disorder and defect, but has not led to any substitution of Ga atoms. Thelocal vibrational mode of Mn substituting the Ga site, calculated by simple mass defect approximation, has been found to be at ~574 cm−1 [28–30] and the absence of any such Raman signal gives corroborative conclusion that MnO2 presence during synthesis has not led to any Mn doping in GaN nanowires. It is to be noted here that the E2(high) mode in GaN is sensitive to stress and the deviation from 569 cm−1 directly gives the measure of the degree of stress . The E2(high) mode shifts to lower wavenumber as MnO2 content is increased. For sample-3, the E2(high) mode appears at ~560 cm−1. This observation indicates that the serrated GaN nanowires might be having a tensile stress. There is a visible shift in A1(LO) mode in GaN nanowires grown in presence of MnO2 compared to that in GaN nanowire grown without MnO2. The mode shifts to 724 cm−1 in sample-2 from that of 730 cm−1 in the case of sample-1. The shift is further increased in sample-3 where the A1(LO) mode peaks at 712 cm−1. The shift to lower wavenumbers is indicative of the fact that the carrier concentration is smaller in serrated GaN nanowires than in the straight GaN nanowires. This is consistent with resistivity measurements which show that the conductivity of serrated nanowires is lower than that of the regular nanowires. The two-probe measurements (Fig. 5) reveal a non-ohmic behavior for both the straight and serrated nanowires; however, the conductivity of the serrated nanowireswas estimated to be up to two orders of magnitude lower than the conductivity of straight nanowires measured under similar conditions. In addition to reduced carrier concentration as discussed above, other factors such as the anisotropic growth direction, enhanced scattering at surfaces may also play a role in reducing the conductivity in the serrated morphology.
At this juncture, we can look at the growth mechanism of GaN nanowires with the presence of Mn vapor in the growth window. At the growth temperature (960°C) inside the quartz tube, the Au catalyst is in liquid phase and the decomposition of solid MnO2 powder gives vaporized Mn, whose concentration is low. As per the phase diagram, the solubility of vaporized Mn in Au rich region is low . Hence, successful doping of Mn into GaN is hindered. Depending on the solubility of the vaporized species, the liquid phase of Au catalyst act as a preferential adsorption site for the vaporized elements until a super saturation stage is reached. The condensation of the vaporized elements then forms an alloy at the liquid-substrate (Au-Si, in the present case) which continues to grow into nanowires. In general, the serrated morphology of GaN is initiated by the liquid Au droplet [17,18]. An oversimplified schematic of the growth process is depicted in Fig. 6.When a semi-continuous Au film (Fig. 6(a), representative SEM image is also displayed) is annealed (in Ar flow of 70 sccm at 960°C for 10 min.) in CVD before GaN growth process, Au droplets of various sizes are formed (Fig. 6(b), SEM image of the annealed Au film is also shown). Depending on the initial effective Au thickness, the shape (and so is the average size) of the droplet varies from spherical to hemisphere. The initial orientation of nanowire growth is dependent on interfacial angle of catalyst Au droplet, i.e, an acute or obtuse contact angle will initiate the outward or inward growth, respectively (right panel of Fig. 6(c)). In principle, when the droplet makes an obtuse angle of contact with the interface, it forces the nanowire to grow inward in order to minimize the free energy. This process leads to compression of Au droplet and the growth process then gradually increases the contact angle (of Au droplet with interface) to 90°, anequilibrium point. Further growth makes the contact angle acute and this initiates the outward growth of nanowires and consequently the Au droplet gets stretched. Hence, the droplet periodically get stretched and compressed whenever the nanowire widens or narrows and this result in serrated feature (right schematic of Fig. 6(d)). However, when the Au droplet size is very small, the initial contact angle could be ~90° (schematic shown on the left side of Fig. 6(c)) and normal growth (straight) proceeds (left schematic of Fig. 6(d)). It is instructive at this stage to recall Fig. 1(b) which shows the simultaneous existence of straight as well as serrated side wall nanowires. The section of Au droplets with larger size, as already discussed, mostly promotes serrated morphology. Therefore, in the presence of MnO2, the higher yield of serrated growth under similar experimental conditions might be arising from those sections of liquid Au whose size is small. We believe that the Mn vapor pressure might exert a force sufficient enough on small Au droplets to modulate the contact angle to off equilibrium angle. In other words, as shown schematically in Fig. 6(e), the small Au droplets with contact angle ~90° could be deformed to hemispheres by the constant pressure exerted by the Mn vapor (indicated by arrows), resulting in obtuse angle of contact with interface. Due to this, possibility of serrated morphology increases. As was confirmed from the Raman spectra, the vaporized Mn introduced a tensile stress into the lattice. This in turn may further facilitate the vertical growth of serrated nanowires. It was in fact seen that the GaN nanowires grown in presence of MnO2 were longer and straighter than that of as-grown GaN nanowires. Higher concentration of vaporized Mn thus seems to enhance the quality as well as quantity of serrated nanowires. It is worth mentioning here that the outer diameter of the serrated GaN nanowire in the present work is approximately that of the initial Au particle size. Hence, initial ratio of the precursor material and effective Au film thickness are the two important parameters that can tailor the morphology, shape and size of the as-grown serrated GaN nanowires.
Periodically serrated GaN nanowires have been grown by chemical vapor deposition method using Au as catalyst. Owing to the low solubility of vaporized Mn in liquid Au, successful doping of Mn could not be achieved in the present growth mode. However, the presence of Mn vapor does exert a force that modulates the interfacial balance at the vapor-liquid-nanowire interface resulting in the modification of the morphology of, otherwise expected straight GaN nanowires. In addition, the Mn vapor introduces tensile strain in the lattice and further promotes the vertical growth of GaN nanowires. Hence an enhanced production of good quality serrated GaN nanowire growth has been successfully achieved. The geometrical aspect of serrated GaN could be easily controlled by initial effective Au film thickness, Ga/N ratio and MnO2 concentration. From PL, micro-Raman, XRD, SEM and conductivity studies we propose a model for formation of serrated nanowires. The GaN nanowires exhibit broad yellow-green luminescence band in addition to red luminescence. The serrated GaN nanowire has larger surface to volume ratio and may have potential sensor applications. Moreover, diameter control along the nanowire length could also be used to enhance device performance.
We gratefully thank Mr. Ajay Jha (TIFR, Mumbai) for assistance with Raman measurements. Zheng Ma and Latika Menon acknowledge NSF ECCS grant #0925285.
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