Facet-driven formation of axial and radial In(Ga)As clusters in GaAs nanowires

Embedding quantum dots in nanowires (NWs) constitutes one promising building block for quantum photonic technologies. Earlier attempts to grow InAs quantum dots on GaAs nanowires were based on the Stranski–Krastanov growth mechanism. Here, we propose a novel strain-driven mechanism to form 3D In-rich clusters on the NW sidewalls and also on the NW top facets. The focus is on ternary InGaAs nanowire quantum dots which are particularly attractive for producing single photons at telecommunication wavelengths. In(Ga)As clusters were realized on the inclined top facets and also on the {11-2} corner facets of GaAs NW arrays by depositing InAs at a high growth temperature (630 °C). High-angle annular dark-field scanning transmission electron microscopy combined with energy-dispersive x-ray spectroscopy confirms that the observed 3D clusters are indeed In-rich. The optical functionality of the as-grown samples was verified using optical technique of cathodoluminescence. Emission maps close to the NW tip shows the presence of optically active emission centers along the NW sidewalls. Our work illustrates how facets can be used to engineer the growth of localized emitters in semiconducting NWs.


Introduction
Single photon emitters (SPEs) are envisioned to play a central role in the future of quantum photonics. Their development is 4 These authors contributed equally to this work.
Original content from this work may be used under the terms of the Creative Commons Attribution 4.0 licence. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. important for applications in quantum computation, quantum communication and quantum cryptography [1]. An ideal single-photon source is identified by a high fidelity and a high degree of anti-bunching (g 2 (0) ∼ 0), high brightness and narrow emission linewidth [2]. In recent years, bright single-photon emitters have been demonstrated by embedding semiconducting quantum dots (QDs) within nanowires (NWs) [3][4][5][6]. Semiconductor QDs are three-dimensionally confined nanostructures with discrete, atom-like emission lines. Due to quantum confinement, electronic levels in a QD can be tuned by changing their size [7] and composition [6]. Conventional QDs are grown by planar epitaxy and embedded in a three-dimensional matrix of a higher bandgap material [8]. However, the total internal reflection at the bulk semiconductor/vacuum interface and the far-field divergence of the emission make the photon extraction highly inefficient from these structures [2,4,9]. NWs, on the other hand, can be designed to operate as photonic waveguides to enhance light extraction. Proper positioning of a QD on the NW axis allows the coupling of QD emission to the fundamental HE 11 mode of the NW waveguide and coupling to other radiation modes is suppressed [4,5,10].
Apart from efficient light extraction, the size and morphology of the NWs offer several degrees of freedom to form low dimensional heterostructures like QDs and quantum wells (QWs). It is possible to obtain axial and radial heterostructures in NWs by varying the material composition along the axial and radial directions, respectively [11,12]. Moreover, NWs provide a wider choice of material combinations with a higher lattice mismatch for forming axial and radial heterostructures. For instance, in the case of III-V semiconductors, one can either change the group-III element or the group-V element to form heterostructures. Some further combinations are also possible with ternary alloys consisting of an additional group-III or group-V element (e.g. InGaAs, InAsP). The large aspect ratio of NWs helps to relax the misfit strain induced by the lattice mismatch on the sidewalls and prevents the occurrence of defects [13]. Out of the available material combinations, InGaAs nanowire quantum dots (NWQDs) are particularly of interest due to their ability to emit photons in the telecommunication wavelength range and with purity in excess of 99% (g 2 (t) < 0.01) [9,14]. Therefore, positioning InGaAs QDs on the axis of GaAs NWs is an attractive choice to realize singlephoton sources.
Most reports on the III-V QDs in NWs are based on NWs grown using gold as the catalyst. However, the incorporation of gold into the grown semiconductor is found to be detrimental for its optical properties [15]. Although self-catalysed growth of defect-free ordered GaAs NW arrays are widely reported, NW axial heterostructures based on this method is still missing. In the case of III-V semiconductors, axial heterostructures are obtained by either changing the group-III element, like GaAs/InAs, GaP/InP, or the group-V element like GaAs/GaSb. More combinations are available if either A or B is a ternary alloy. GaAs QDs in AlGaAs NWs, InAsP QDs in InP NWs, and InGaAs QDs in GaAs NWs are a few examples. Despite the availability of a wide variety of material combinations, challenges like the reservoir effect [16] and kinking [17] limit their successful implementation.
Here, we take an alternative approach for switching between two materials (InAs and GaAs). The Ga droplet, used as a growth catalyst, is completely consumed before the deposition of InAs. Formation of axial and radial localized InGaAs clusters within self-catalysed GaAs NWs is observed. Specific focus is given to the structural characterization of the obtained heterostructures using transmission electron microscopy (TEM). Our findings provide new insights into how facets can be used to engineer the growth of QDs in semiconductor NWs.

Experimental details
The GaAs NWs studied here were grown by the self-catalysed vapour-liquid-solid (VLS) method in a DCA P600 molecular beam epitaxy (MBE) system. A Si (111) wafer coated with an SiO 2 mask was patterned by e-beam lithography and used as the growth substrate. The NW growth was carried out at a substrate temperature of 630 • C and at Ga and As 4 beam equivalent pressures (BEPs) of 1.4 × 10 −7 Torr and 2 × 10 −6 Torr, respectively. The Ga supply was interrupted after an hour of NW growth. The Ga droplet used as a catalyst in the NW growth was then completely consumed under an As 4 flux of 7.2 × 10 −7 Torr. InAs was deposited for 5 s at a substrate temperature of 630 • C under an As 4 BEP of 5.6 × 10 −6 Torr and at a V/III ratio of 45. The morphology of the grown NW arrays was examined in a Zeiss Merlin scanning electron microscope (SEM). For structural investigation, the as-grown NWs were transferred onto a carbon-coated Cu grids and studied in a FEI Talos TEM operated at 200 keV in high-resolution and high-angle annular dark-field scanning TEM (HAADF-STEM) modes. Elemental maps were obtained by energydispersive x-ray (EDX) spectroscopy. The optical properties of the NWs were studied by micro-photoluminescence (µ-PL) in a cryostat operating at around 4 K under an optical excitation provided by the 632.8 nm line of He-Ne laser. The PL signal was collected into a 300 mm focal length spectrometer and dispersed by a 150 l mm −1 grating onto a Peltier-cooled CCD. Cathodoluminescence (CL) measurements on the NWs were carried out in an Attolight Rosa 4634 CL-SEM microscope. The measurements were carried out at 10 K with a beam energy of 5 keV and probe currents of 5-20 nA. The CL signal was collected into a 320 mm focal length spectrometer and dispersed by a 150 l mm −1 grating onto an IR-sensitive Peltiercooled InGaAs PDA detector. A selected-area electron diffraction (SAED) pattern from the encircled regions in figures 1(b) and (d) is provided as an inset which reveals the crystal structure of these sections of the NW. All the NWs showed a predominantly zincblende (ZB) phase at the bottom (figure 1(b)), followed by a transition region composed of a dense region of stacking faults (figure 1 (c)). Following the transition region, the upper segment of the NWs is pure wurtzite (WZ) as shown in figure 1(d). The observed transitions in the crystal phase are due to the variation in the Ga droplet contact angle [18][19][20]. Variations in the droplet contact angle modify the capillary forces at the triple phase line (TPL) and control the formation of ZB and WZ phases [21]. Recent in-situ TEM growth experiments have shown that ZB phase is favoured at small (<100 • ) and large (>125 • ) contact angles while WZ phase is favoured at any intermediate contact angle [22].

Results and discussion
In the VLS growth of GaAs NWs, the Ga droplet acts as a non-stationary reservoir of Ga [23,24]. The relative Ga flux during growth determines the droplet size and the contact angle. An As-rich growth condition shrinks the droplet and reduces the contact angle, whereas, a Ga-rich condition swells the droplet and increases the contact angle. Therefore, a predominantly ZB crystal structure at the bottom of the NW is due to Ga-rich conditions at the beginning of the growth. After reaching the desired length, the Ga flux is terminated, and the droplet is consumed under a constant As 4 flux. As the droplet shrinks, the contact angle also decreases gradually, creating a transition region full of stacking faults. When the contact angle falls below 125 • , the WZ phase becomes favourable, and the consumption of the remaining Ga droplet results in a pure WZ phase. Finally, at a very small droplet size, the NW tip is again ZB but has a different facet morphology (figure 2). For a core-shell NW heterostructure, defects in the NW core template into the shell [25]. Thus, GaAs NW core with a defectfree tip is essential for the growth of axial heterostructures. The SEM images of the NW tip after Ga droplet consumption at 630 • C are shown in figures 2(a)-(e). Figures 2(f)-(h) show high-resolution TEM images of the NW tip taken along the <1-10> zone axis. The tilted facets are indexed by measuring the angles made with the <-1-1-1> NW growth plane. The tip of the NW showed a pure ZB phase and is free of stacking faults (figures 2(g) and (h)). However, non-uniformities in the NW tip morphology are observed. A major fraction of the NW tips observed in SEM showed a truncated pyramidal morphology with a flat triangular {111}B facet at the tip as evident in figures 2(b)-(d). The remaining fraction of NWs showed a completely pinched-off NW tip as shown in figure 2(e). In crystal growth, facet formation is governed by the slowest growing plane. Any variation in the growth conditions (like V/III ratio or temperature) affects the growth rate of different planes and dictates the final facet morphology. Under the conditions studied here, we believe the group-III and group-V reevaporation rates to be substantial and varying from one NW to another. Therefore, the observed non-uniformities in the tip morphology can be attributed to the local variation in growth conditions.

Formation of axial In(Ga)As clusters:
Having obtained a good understanding of the GaAs NW tip morphology and crystal structure, the effect of InAs deposition on GaAs NWs was investigated. STEM measurements were performed to study the preferential accumulation of InAs on the NWs. Figure 3(a) shows the HAADF-STEM image of the upper region of the NW after InAs deposition. Due to the higher atomic number (Z) of In with respect to Ga, we attribute the regions with a brighter intensity to the presence of In(Ga)As clusters. These clusters are observed on the top facet (axial clusters) as well as along the NW side facets (radial clusters). Figure 3(c) shows a magnified HAADF image of the GaAs NW tip after InAs deposition as indicated by a red rectangular region in figure 3(a). The corresponding In and Ga EDX maps are shown in figures 3(d) and (e). In(Ga)As clusters accumulate at the tip of the NW whereas, pure InAs accumulation is observed on both the inclined top facets. When the amount of InAs deposited is reduced, the InAs clusters become even smaller and are observed only on one of the inclined facets at the NW tip (supplementary information (https://stacks.iop.org/JOPT/22/084002/mmedia)). We   clearly evident from SEM images in figures 2(b)-(e). The InAs clusters accumulating on the corner facets at the tip are similar to the clusters nucleating at the corner facets on the NW sidewalls, as discussed in section 3.2. Further investigation is needed to determine the exact location of these clusters. This, however, shows that engineering of the NW top facet could help in obtaining axial InAs clusters.
The Stranski-Krastonov growth of three-dimensional InAs islands on these surfaces has been realized only after covering the facets with a thin layer of surfactant material [28,31]. In our case, we observe the formation of three-dimensional In-rich clusters on GaAs NW side facets by depositing InAs layer at a higher substrate temperature. In HAADF images of figures 3(a) and (b), which are taken along top and bottom parts of the NW, we observe regions with brighter contrast all along the NW sidewalls. This indicates the presence of radial In(Ga)As clusters along with axial clusters on GaAs NWs. This is also evident from the EDX elemental map of In overlapped with HAADF images of NW sidewalls in figures 3(f) and (g). The location of the observed In-rich radial clusters in figure 3(f) clearly suggests that they are nucleating on the edges separating the two {110} NW side facets.
In the following discussion, we propose a material transport model to explain the temperature-dependent formation of 3-D In(Ga)As clusters. We base our proposition on strain energy minimization. Strain energy minimization [32] and gradient in chemical potential [26] are the two known driving forces for adatom diffusion between different NW facets. The convex nature of the {11-2} corner facets makes them favourable sites for strain relaxation and the nucleation of 3-D structures [33]. At a higher substrate temperature, adatoms have enough kinetic energy to overcome the chemical potential barrier. Thus, material transport due to strain relaxation becomes dominant. This, in turn, promotes preferential accumulation of the impinging adatoms at the corner facets. This results in the formation of In-rich clusters along the {11-2} corner facets of the NW. A schematic of the proposed mechanism is shown in figure 4(c). At higher temperatures, In(Ga)As clusters preferentially accumulate at the corner facets indicated by red arrows, as observed in figure 3(f). In contrast, growth at a lower substrate temperature in the absence of a surfactant leads to the formation of a thin 2-D In(Ga)As shell [34]. Figures 4(a) and (b) show SEM images of upper and lower regions, respectively, of a NW sample grown at a lower substrate temperature (440 • C). The InAs deposition was carried out for 5 mins under an As 4 BEP of 5.6 × 10 −6 Torr and a V/III ratio of 45. The formation of a thin discontinuous 2-D shell is observed. The lower diffusivity of In adatoms at this growth temperature leads to a high density of such clusters at the bottom of the NW ( figure 4(b)) whereas a lower density is seen in the upper part of the NW ( figure 4(a)).

Optical characterization of the In(Ga)As clusters:
To investigate the optical properties, CL and PL measurements were carried out on NW structures similar to the ones studied by STEM-EDX. However, to reduce the size of In(Ga)As clusters, InAs was deposited for 3 s under an In BEP of 2.2 × 10 −8 Torr. The NWs were subsequently capped with a GaAs shell for confinement. The GaAs shell growth was carried out at a substrate temperature of 440 • C and under a high V/III ratio of 70 to ensure growth of a conformal 2-D GaAs shell. Ga and As 4 BEPs were maintained at 1.4 × 10 −7 and 1.0 × 10 −5 Torr, respectively. Figure 4(d) shows the HAADF  image of a NW after shell capping. The tip of the NW showed a high density of defects upon capping with a GaAs shell. We also note that the HAADF contrast as seen in figure 3 from the In-rich clusters is not visible upon shell capping. Figure 5 shows the CL spectra of two representative NWs from the sample after shell capping. The CL spectra correspond to the emission of a region close to the NW tip and are averaged through the region indicated in the SEM images in the insets. The insets on the left and the right show the energyfiltered CL maps and the SEM image of the NWs, respectively. The colour coding is mentioned on the respective energyfiltered CL maps in the insets. The spectra in figures 5(a) and (b) both show multiple peaks originating from different emission zones. Some emission zones appear as a strip across the whole NW diameter (1.207 eV) while others extend over very few pixels (1.225 eV, 1.24 eV, 1.271 eV, 1.291 eV). In agreement with the structural characterization results (figure 3), all the emission zones observed in the CL maps are in close proximity to each other. The emission zones in the CL maps of figure 5(a) are in agreement with the location of In(Ga)As clusters seen in figures 3(f) and (g). Due to the high density of defects at the tip, no emission is observed from this region of the NWs. In figure 5(b), the emission zones are near the stacking faults. This is evident from the emission map corresponding to 1.225 eV which indicates accumulation of In-rich clusters in the region with a high density of stacking faults. This is in agreement with the HAADF image in the inset of figure 5(b) where regions with brighter contrast are seen in the region full of stacking faults.
Due to the carrier diffusion length and excitation volume, the exact location of each localized clusters is challenging to resolve. However, the occurrence of multiple emission peaks suggests the presence of In-rich clusters all along the NW sidewalls. To confirm the presence of very thin In-rich clusters along the NW sidewalls, position-dependent PL measurements were carried out as shown in figure 6. The PL spectra were acquired from a single NW lying horizontally on the as-grown substrate. The background emission from the underlying substrate was absent and ensures that the spectra shown here indeed originate purely from the NW heterostructure. The spectra acquired at very low excitation powers reveal sharp QD-like emission lines. The narrow peaks with linewidths in the sub-meV range indicate that the emission originates from the In(Ga)As clusters confined on the NW sidewalls.

Conclusion
In summary, we have studied the effect of NW growth termination on the morphology of the GaAs NW tips obtained in ordered arrays. Depending on the NW morphology and growth temperature we have demonstrated the formation of axial and radial In(Ga)As clusters. GaAs NW tips showed a truncated pyramidal morphology with a flat triangular {111}B facet at the top. Moreover, the tips are found to be free of stacking faults and have a pure ZB crystal structure. In addition, In(Ga)As clusters were obtained all along the corner facets of the NW sidewalls. The occurrence of radial In(Ga)As clusters was confirmed by scanning TEM and low-temperature CL and PL measurements. We believe that the In(Ga)As cluster formation is driven by the convex nature of the {11-2} corner facets. This approach of facet-driven nanostructure self-assembly can be engineered to grow axial and radial QDs in latticemismatched NW heterostructures. In addition, they could also be exploited to make novel quantum structures ranging from 3-D islands to zig-zag shaped nano-rings on NWs [29]. The QD-in-nanowires could potentially contribute to single photon emitter devices with large repetition rates owing to small radiative decay lifetimes [35]. Considering the shortest excitonic lifetime measured in QD-in-nanowire systems (0.42 ns) [36], repetition rates could in principle be as high as 2 GHz. However, this high emission rate is only achievable by optical excitation. Challenges currently reside in the electrical excitation of QD-in-nanowires. Nonetheless, with the availability of sophisticated manipulation and processing techniques, the integration of nanowire QD emitters into silicon photonics operating at telecom wavelengths approaches the possible [37].