Buk, Carmen, Dimitris Andrea S Giacomo

Martina Juan

Nicolas

Quentin Nicoletta

Vivek Thanks

First of all, I would like to thank the members of the jury for accepting to be in my PhD defence. In particular, I would like

Table of contents

In modern electronics and optoelectronics, fabrication of high quality monocrystalline layers (called 'active layers') is a challenge. This problem is more pronounced for gallium nitride material growth as the native substrate is of very high cost and the hetero-epitaxial substrates have large lattice mismatch and thermal expansion coefficient mismatch. As a result, GaN layer is grown mostly on sapphire and silicon substrates with buffer layers, to improve the layer quality. This increases the cost and complexity of the material growth and the active layer is not completely defect-free.

A major advantage of semiconductor nanowires (NWs) is the possibility to integrate these nanomaterials on various substrates: large lattice mismatches with the substrate can be accommodated with no generation of extended defects in these nanostructures and the difference of thermal expansion coefficients is less critical than for two-dimensional layer epitaxy. These perspectives are particularly attractive for the IIInitrides, considering the fact that the epitaxial growth of these compounds suffers from a lack of an ideal substrate.

Fabrication of NW based devices requires vertical growth of NWs and conventionally it necessitates a crystalline substrate which is generally a 400-600 μm thick wafer, sliced from a bulk single crystal. The substrate is generally much thicker than the height of the NWs which represent only a few μm or even less. In many cases, the substrate has no further crucial role after the epitaxial growth. In certain situations, it is even beneficial to remove it and to transfer the NWs on another support presenting properties more adapted to the device operation (thermal management, transparency, electrical conductivity). Such a process is delicate, costly and usually the substrate is lost.

We wanted to develop an entirely novel technology which eliminates the use of thick mono-crystalline wafers for epitaxy. In this context, we examined the use of novel templates for growing GaN NWs of high structural quality by plasma assisted molecular beam epitaxy (PAMBE). We explore three approaches with a common feature: the initial support is an amorphous substrate which could be a cost-effective solution for developing future photonic devices. Then, the idea is to use a very thin crystalline material deposited on the bulk amorphous substrate to promote the epitaxial growth of GaN NWs.

With this idea, we explored in detail the following templates:

Organization of this thesis

Chapter 1 provides a general introduction to the history of gallium nitride crystal growth and the problem of having a suitable substrate for epitaxy of GaN. It also gives a general introduction about NWs by describing advantages of the NWs over the thin-film growth, the substrate requirement and the basic mechanisms involved in NW growth using plasma assisted molecular beam epitaxy (PAMBE). The motivation for this project is explained in detail and the state-of-the-art results for NW growth on non-conventional substrates is presented. In this chapter the overall objective of this project is also set.

Chapter 2 comprises of introduction and state-of-the-art results for growth of GaN NWs on Si(111). We have used the state-of-the-art growth conditions to fabricate NWs on Si(111) which will serve as our reference sample. The growth procedure we used to obtain GaN NWs on Si(111) are explained in detail. This chapter also has the steps involved in calibration of growth flux and temperature and a brief introduction to the characterisation techniques we used such as x-ray diffraction measurements (XRD) and photoluminescence (PL) measurements, which are common sections for the following chapters.

Chapter 3 provides introduction to the concept of using thin-silicon-layers, prepared by Aluminium induced crystallisation (AIC-Si), as substrate for GaN NWs. I have detailed the mechanism involved in formation of AIC-Si and the means of obtaining the same using sputtering. State-of-the-art results for using AIC-Si in NW growth is presented in this chapter. Using standard growth conditions, NWs are synthesised on AIC-Si layers by self-induced growth and characterised using different techniques like XRD, PL and energy dispersive x-ray spectroscopy (EDS). Selective area growth of NWs on the AIC-Si layers is demonstrated by patterning the AIC-Si layer to produce 'nano-substrates'.

Chapter 4 details usage of silica wafers as the substrate for GaN NW growth. We explored thermal silica and fused silica for direct growth of GaN NWs on them using MBE. A brief state-of-the-art results for NW growth on amorphous substrates are presented. NWs are characterised using different techniques such as XRD, transmission electron microscopy (TEM), KOH etching, EDS and PL and compared to the NWs obtained on Si(111).

Chapter 5 explains GaN NW growth on graphene layers. In this chapter, I have provided a brief introduction about van der Waals epitaxy and understanding of the same found in literature. Introduction to the different graphene layers used and the detailed steps to prepare these layers are provided. Morphology of the NWs grown is analysed and other characterisation techniques such as XRD, TEM and PL are used to compare the NW qualitatively with our reference sample. The epitaxial relationship of GaN with graphene is studied in detail using two different methods: Scanning electron microscopy (SEM) and TEM. A strain sharing model is illustrated which explains the NW growth behaviour with respect to the number of monolayers of graphene. Finally, proof of using e-beam patterned graphene patches for selective area growth of NW arrays is provided.

Chapter 1

Necessity for a novel substrate for gallium nitride nanowires

1.1 Gallium Nitride: An introduction

A brief history of GaN crystal growth

In 1931, first growth of polycrystalline Gallium nitride (GaN) was obtained by reacting metallic gallium with ammonia gas at 900-1000 o C 1 . In the 1950's-60's, light emitting diodes (LEDs) with different wavelengths from infra-red to green were available, except for the blue colour 2 . At the same period, it was demonstrated that it is possible to fabricate LEDs with GaN 3 . However, it was tough to get good quality single-crystalline GaN.

By 1969, single crystalline GaN with a good crystal quality was grown using vapour phase growth technique. This semiconductor was reported to have a direct bandgap of 3.39eV 4 . This looked promising to obtain blue LEDs using GaN and thereby combining with other already existing LED colours, it might be possible to fabricate lighting devices completely based on solid-state materials.

In 1971, GaN on Al2O3 and SiC substrates using MOVPE was demonstrated 5 and GaN LED using i-n doped junctions was demonstrated by Pankove et al. 6 . However, it was difficult to improve the crystal quality and also to obtain p-doped GaN which is crucial to obtain high efficiency LEDs using p-n junctions. As a result, after a few years of intense research, many leading groups stopped researching on GaN and moved to research on other materials to obtain blue LEDs in the late 1970s 7 . At this time, a few started to work on GaN growth by MBE and MOVPE. It was in 1986 where the major breakthrough came to obtain a better quality GaN crystal. Amano et al.

reported high quality GaN, grown with AlN buffer layer on sapphire substrate using MOCVD technique 8 .

After detailed studies on using AlN buffer layers 9,10 , the same group a couple of years later, came up with results for p-doped GaN, showing an improvement in p-type conduction by irradiating doped films with low energy electrons and p-n junction LEDs were fabricated for the first time using this AlN buffer layer 11 . In the meantime, two-flow MOCVD system was developed by Nakamura et al. 12 to have a good control over the reactant gas flow and thus to improve the uniformity of GaN layers. It was shown that high quality GaN can be grown with a low temperature GaN buffer 13 and GaN p-n junction blue LEDs were demonstrated for the first time 14 . Immediately after, Nakamura showed that it is possible to grow high quality p-doped GaN by just annealing the doped GaN layers 15 . Since then, the research on GaN, in terms of publication, has been on rise until now 16 . These above mentioned crucial works of Amano, Akasaki and Nakamura led to their Nobel prize in 2014.

Though it's the LED market that pushed GaN based research up in the initial years, the material based heterostructures are being broadly explored for other applications such as high-power electronics due to its high breakdown voltage 17 . Due to its high electron mobility (~400 cm 2 /Vs), GaN is explored for high frequency applications since as early as 1994 18 .

Other than these applications, it is explored for photovoltaics as indium gallium nitride (InGaN) has the capability to cover almost the entire solar spectrum, from 0.65 eV to 3.42 eV, though the reported efficiency is low due to high number of defects introduced by the high incorporation of indium in InGaN films 19 .

Substrates for epitaxial growth of GaN

Due to the lack of cheap native substrates (GaN), through all these years in the research of GaN thin film growth, one of the important issues to be solved in advancement of GaN-based technology is improving the material quality. As GaN is not present as a natural mineral (like Si) and due to the high melting temperature, conventional bulk substrate preparation methods such as Czochralski or Bridgeman techniques cannot be used.

As a result, GaN has to be grown from gallium and nitrogen precursors to obtain single crystalline bulks. These days GaN substrates are grown by ammonothermal method and hydride vapour phase epitaxy (HVPE), which makes the substrate limited in size and much more expensive than the conventional substrates. Hence in most of the cases for GaN based device fabrication, GaN is grown on foreign substrates which are having high lattice mismatch and high thermal expansion coefficient mismatch. Fig. 3 shows the lattice mismatch (with GaN) versus cost for different heteroepitaxial substrates available in market for GaN growth. Sapphire, silicon carbide (SiC) and silicon (Si) are the most commonly used substrates for GaN growth with sapphire dominating the market with a share of approximately 80%.

Figure 2: X-ray rocking curve showing the (0006) diffraction peaks for GaN layers grown with AlN buffer (dotted line) and without AlN buffer (continuous line). Taken and edited from Amano et al. 8 As the substrate is a quasi-infinite material compared to the small amount of material grown on it, most of the misfit strain must be accommodated in the epilayer. As strain energy accumulates in the growing film, defects, dislocation and cracks are generated to recover the bulk lattice parameters. This is detrimental for device performances as these defects lead to undesired effects like non-radiative recombination, scattering centres and reduced breakdown voltage.

The possibly induced defects can be classified into three categories: point defects, linear defects, planar defects.

Point defects (anti-sites, interstitials, vacancies) can arise at any part of the material grown as they are based on the local arrangement of particular atoms. The other two types of defects, linear defects and planar defects are extended defects which are mainly associated to the mismatch between the epitaxial materials and the foreign substrates. Wu et al., studied in detail the dislocations observed in GaN films and showed that approximately 70% of dislocations observed are pure-edge dislocations 20 which are, in another work, shown to generate during the initial stages of growth and propagate in the growth direction to the surface of the films 21,22 .

Other types such as Shockley and Frank partial dislocations are also observed close to the film-substrate interface. Apart from these dislocations, the stress induced by the mismatch is released in terms of cracks and stacking faults. Inversion domain boundaries are also observed to be formed during the initial stages of growth from the film/substrate interface 20 . Thus most of the defects observed in GaN films with high defect density (10 8 -10 10 cm -2 ) originate at the film/substrate interface. As of now in industry, the quality of GaN material is optimized by use of intellectual-property (IP) protected buffer layers. This again increases the cost of manufacturing due to the additional technological steps which, nonetheless, do not lead to completely defectfree thin films 23,24,25 .

Introduction to Nanowires

Advantages of nanowires

The problem of substrate is more pronounced for growth of thin films on lattice-mismatched surfaces. This is mainly due to the fact that the defects originating at the film/substrate interface due to the misfit of lattices, cannot propagate to the quasi infinite width of the films and hence disseminates all the way to the surface of the film. This is different in the case of nanostructures with reduced lateral dimensions. It has been shown that plastic relaxation occurs in planar thin films whereas elastic relaxation happens when the deposited material forms three-dimensional islands 26 . Nanowires (NWs) are quasi-one dimensional materials with diameters in the order of tens of nms and with very high aspect ratios (~ >10). This gives them an edge over the thin films by the fact that misfit strain can be accommodated by elastic relaxation at the lateral free-surface. In fact, F. Glas showed that for a certain % of misfit between the substrate and the NW material, there is a critical diameter of the NW below which there should not be any dislocation introduced in them, i.e., the NWs should be able to deform elastically to fully accommodate the misfit stress 27 .

Critical dimensions for misfit of up to 7% were calculated and it can be seen that 20 nm is the maximum NW radius that could accommodate strain elastically for this value of misfit. Moreover, if the NW diameter is above this critical diameter but remains relatively thin, the dislocations created at the base of the NW can easily reach the lateral free-surface and do not propagate to the top. When we consider a thinfilm as a structure of infinite radius (compared to NWs), it can be understood that much higher quality material can be obtained in NWs for the same amount of misfit, proving NWs are one of the most viable solutions to get high quality GaN for device fabrication.

Apart from providing a solution to the substrate issue, NWs also have the advantage of being bottomup grown building blocks from which devices can be fabricated without processing steps on the active device area, like reactive ion etching, which could damage the material. Lithography can be used to pattern the substrate before growth, providing predefined the filling factor (pitch of a regular array and NW diameter), the absorption of NWs can be enhanced to provide better absorption than thin films in all the wavelengths 28 . There are different growth methods available to fabricate GaN nanowires, however the two main methods are metal organic chemical vapour deposition (MOCVD) and molecular beam epitaxy (MBE). MBE is particularly interesting due to the possibility of obtaining high quality material at a comparatively lower growth temperature. In MBE, GaN NW growth can be obtained by two methods, catalyst-assisted and catalyst-free.

NW growth by PAMBE

Catalyst-assisted method is known to induce contamination and defects in the NWs and for this reason, catalystfree method is mostly preferred 29 .

A detailed picture of the MBE chamber is given in annex A.3. Growth is carried out in ultra-high vacuum (UHV), the background pressure being typically 10 -10 Torr, with the substrate kept at an elevated temperature.

The Once the stable nuclei are formed, surface energies, strain and kinetics determine the type of structures being grown with basis of three possible growth modes:

1. Frank-van der Merwe (FM) mode: It is also called layer-by-layer growth method, where the nucleation of a second layer is possible only after the completion of the previous layer. This is mainly due to the stronger interaction of atoms to the surface rather than to the neighbouring atoms.

2. Volmer-Weber (VW) mode: It is the opposite to FM growth mode by the fact that the interaction with the neighbouring atoms are higher than that with the surface. As a result, the structures grown are in the form of islands, because of which, it is also known as 3D growth mode. formation of defects (Fig. 8). As a result, like in the thin-film growth, bulk crystalline substrates are used for NW growth. However, this approach results in large material loss and is cost inefficient (compared for instance to the use of cheaper amorphous substrates). Moreover, NWs are less demanding with respect to the substrate crystallinity. Indeed, their small footprint on the surface and the spacing between them can accommodate to a polycrystalline surface to some extent. As will be seen in Chapter 3, a polycrystalline surface with grain size typically greater than the NW diameter and with a high fiber texture to preserve their oriented growth, is suitable to obtain vertical NWs (Fig. 9). Alternatively, we will show in Chapter 4 that, under proper conditions, GaN NWs can grow vertically even on an amorphous surface. This extends considerably the choice of possible supports which can be selected according to the targeted application.

generally

In addition, the recent development of flexible optoelectronic devices leads also to reconsider the bulk monocrystalline substrate approach. In this context, NW ensembles as active epitaxial material are good candidates because they can withstand strong deformation of the support without degradation. This characteristic is illustrated by recent fabrications of flexible devices using NWs, where after the growth on bulk-crystalline substrate, the NWs are transferred onto a flexible support 33,34 . Considering that the main requirement for substrates to be used for epitaxial growth is that their surface should be crystalline, the ultimate choice lies in two-dimensional (2D) materials. Among them, graphene is the most popular and we will use it in Chapter 5 for GaN NW growth.

Nowadays a few research groups are focussing on obtaining an alternative substrate for NW growth. In Taken from Ikejiri et al. 36 Choi et al. have succeeded to grow good quality GaN nanostructures (pyramids) on glass supports by using thin-Ti layer as the orientation layer 37 . Sputtered thin-Ti layers are said to have 'fiber-texture', which is the feature of having same orientation along one common axis though the in-plane orientation is random (explained in detail in section 3.1.3.1 of chapter 3). On top of this orientation layer, a GaN buffer layer was grown to act as a nucleation layer for further growth. Mask material was deposited on top of the nucleation layer and openings were done in the mask to grow GaN crystallites. LEDs were fabricated using deposition of a transparent and conductive oxide layer as the top contact and the underlying Ti layer as the bottom contact, thus utilising the orientation layer post-growth as well.

Other state-of-the-art results for non-conventional substrates which include metal induced crystallised layer, glass and graphene are presented in the respective chapters.

Objective of this project

The main goal of this project is to identify a suitable substrate for GaN NW growth by plasma assisted molecular beam epitaxy. A potential polycrystalline substrate should mainly possess the following two features:

-A fiber-texture necessary to get definite orientation of the NWs. In principle the fiber-texture should be parallel to the normal of the substrate in order to get vertical NWs.

-Large grains compared to the footprint of the NWs in order to reduce the probability of NWs nucleating at the grain boundaries which could lead to defects in NWs. The substrates will also be assessed based on the quality of NWs obtained on them. The NWs will be examined for the crystal quality and optical quality. In principal, three substrates will be assessed in detail for growth of GaN NWs by PAMBE:

-Aluminium-induced-crystallised Si (AIC-Si): Unlike random orientations of polycrystalline Si prepared by annealing of a-Si, AIC-Si could be prepared with a high fiber texture necessary for vertical growth of NWs.

-Glass (fused silica): In the progress of the project, we found fused silica to be of interest though it does not possess any sort of crystalline order necessary to obtain vertical growth. In Chapter 4, it will be shown that the roughness of such an amorphous surface plays a major role in obtaining high vertical growth of GaN NWs.

-Graphene: As being the thinnest crystalline layer possible, graphene is naturally a substrate of interest.

This 2D material is also very little explored for growth of GaN NWs at the moment.

Chapter 2

Conventional gallium nitride nanowires growth on Si(111) bulk substrate

2.1 Introduction and state-of-the-art

Brief introduction

Silicon is one of the most abundant materials available and it is widely used as a substrate for the epitaxy of semiconductor thin films. Due to its mass production and well developed crystallisation process, it is the cheapest monocrystalline bulk wafer material available with a high crystal quality. Naturally, for these reasons, Si is also a substrate of interest for LED fabrication. Si has a diamond-cubic crystal structure with a lattice parameter aSi = 5.431 Å. The {111} planes of Si have an atomic arrangement similar to that of the {0001} planes of III-nitrides in wurtzite phase. This is favourable to an epitaxial relationship between these two materials. The interatomic distances to be compared are the basal lattice parameter of GaN (aGaN = 0.3189 nm)

and aSi/2 = 0.3840 nm. Though the corresponding lattice mismatch is high (17%), attempts of thin-film growth of GaN on Si (111) were made as early as 1971 in the same chamber.

Our reference samples are fabricated on (111) Si substrates with the following features:

-Quarter of a 2-inch wafer with 250 µm thickness -High n-type doping (phosphorus doping)

State-of-the-art

Growth of GaN NWs on Si (111) is being studied in detail by various research groups. Most of the reports for GaN NWs on Si (111) can be broadly classified into two approaches:

-Formation of in-situ SiN layer before GaN NW growth: Based on the conditions at which the SiN layer is formed, the presence of this layer leads to random in-plane orientation 3 and reduced verticality of NWs.

-Formation of in-situ AlN layers before GaN NW growth: In order to avoid the formation of above mentioned SiN layer, AlN is deposited.

In this chapter, we focus only on the growth of GaN NWs on Si (111) with AlN pre-deposition and I would like to present some of the latest reports on this approach.

In most of the studies, this AlN pre-deposition is obtained by exposure of clean Si substrate to Al and then exposing it to N. Auzelle et al. 4 studied AlN pre-deposition by 2 different approaches: Al first approach and N first approach. In both the cases AlN pedestals are found and on GaN NW growth, NWs grew on the pedestals.

On Al first approach, GaN first nucleates as pyramidal structures and later NW nuclei appears and both co-exist after a long growth time. The pyramids coalesce to form 2D layers between the NWs. However, in the N-first approach, NWs also nucleate on the AlN surface.

Influence of AlN layer thickness on the morphology and nucleation time of GaN NWs was studied in detail in the report by Musolino et al. 5 . It was found that 4 nm thick AlN buffer layer enables the growth of wellaligned GaN NWs with a short nucleation time. With a SiOx mask, they have shown that it is possible to obtain selective area growth of the NWs on the AlN buffer. LEDs were fabricated from NW ensembles grown on the AlN buffer layers which were not found to have a detrimental effect on electrical injection. From applications point of view, GaN NWs grown on Si (111) with AlN pre-deposition is being explored for piezoelectric devices as well. From our group, Jamond et al. 7 have demonstrated fabrication of piezoelectric generators based on GaN NWs grown on Si (111) substrates. With a prototype, they showed that power density of 12.7 mW.cm -3 is achievable which is the state-of-the-art result for piezo-power generation using NWs.

Standard growth conditions and results

Substrate preparation

The chemical and physical properties of the substrate surface have a big influence on the phenomenon occurring with adatoms on the surface. Depending on the local environment thereof, chemical nature of the atoms constituting the surface, their crystalline arrangement and the type of site (eg., dangling bonds), the impinging flux will react differently with the surface and the nucleation of structures can be more or less efficient due to the variation in different phenomenon mentioned in section 1.2.2. Hence it is necessary to have a clean surface to obtain controlled and reproducible results. Si substrate when exposed to air, normally forms a thin amorphous oxide layer (called native oxide), with a thickness of about 20 Å 8 and a layer of carbon contamination.

For our samples, cleaning is done in two steps. Cleaning with chemicals is made before the introduction of the sample into the MBE and in-situ cleaning is done to remove the final oxide contaminations from the surface. Steps 1-3 correspond to a cleaning with conventionnal solvents to remove the organic residues from the surface. Following that, steps 4-6 serve for removing the native oxide, then for intentional formation of an oxide lawyer along with the carbon residues using a UV lamp and removal of this intentional oxide by immersing the sample again in a diluted HF bath. The last HF dip allows to passivate the surface with Si-H bonds 9 . This helps to reduce the oxidation when exposed to air while the sample is being transferred from the chemical bench to the MBE introduction chamber.

IN-SITU CLEANING: Immediately after the above mentioned chemical cleaning procedure, the Si quarter is mounted on a molybloc (molybdenum holders) and placed in vacuum at the introduction chamber. The sample is exposed to air for less than one minute before being placed under vacuum. The introduction chamber is then baked overnight for 10 hours to remove any atmospheric moisture introduced during the transfer in ambient atmosphere.

Before the growth, the substrate is heated at 860 °C for 15 minutes to remove the final traces of native SiO2 which could have formed during the exposure to air between the chemical cleaning and the introduction into the MBE. A perfectly clean Si surface is identified by the appearance of 7 by 7 reconstruction during a temperature ramp between 600 °C and 830 °C, which will be seen in detail in the following section.

Substrate temperature calibration

Temperature measurement is an important yet difficult task under ultra-high vacuum (UHV) conditions.

Our MBE is equipped with two means of measuring the temperature.

-Thermocouple: a thermocouple is kept in close proximity to the substrate on its holder.

-Pyrometer: a pyrometer deduces the temperature by measuring the infra-red emission from the sample surface at 900 nm wavelength.

The latter technique is more reliable provided that calibration is carefully done and the change in temperature reading is faster by pyrometer compared to the thermocouple. The thermocouple reading has generally a large offset (for instance, thermocouple reads 700 o C for a surface temperature of approximately 830 o C). However, the pyrometer measurements are also affected by different aspects:

-Angle at which the measurement is performed (emission is viewed) -Cleanliness of the glass window by which the measurements are taken -Change in the emissivity due to the use of different substrate materials -Change in the emissivity and interference effects with respect to the structure grown on the surface In order to overcome many of these effects, we calibrated the pyrometer before every growth using the surface reconstructions appearing on Si(111). This is done by considering the temperature of disappearance of the 7 x 7 reconstruction. In the literature, different temperatures associated with 7 x 7 disappearance can be found. They range from 827 °C to 870 °C. We have considered 830 °C as the temperature for disappearance 10 and shift of pyrometer is calibrated based on this reading between samples. For example, immediately after changing the window, the pyrometer reads 860 °C for disappearance of 7 by 7 RHEED pattern. In this case +30 °C is considered as the offset of temperature and all the temperature readings during this sample growth are based on this value. This calibration procedure is particularly well adapted to our growth experiments which are performed at temperatures close to 830 °C.

Flux calibration

GaN NW growth is usually carried out under nitrogen rich conditions (N/Ga >1). In accordance with the change in flux ratio, the morphology of the NWs change. Thus, in addition to the growth temperature, growth fluxes are also important parameters to be controlled.

Gallium and aluminium fluxes are generated using a Knudsen cell and vary as a function of the temperatures of the cells. The flux is measured by using an ion gauge, which is extendable to reach the front of the substrate holder. Active nitrogen species are produced by means of dry N2 of high purity, fed through a flow controller and by generation of plasma in a boron nitride cavity excited by a radio frequency (RF) source. Excited species (N atoms, excited N2, neutral and ionised) escape from the cavity in the substrate direction via small holes. As a result, N flux is controlled in terms of the SCCM (standard cubic centimetres per minute) flow and RF power of the plasma source. However, the ion gauge does not give a reliable measurement of N flux, mainly because of two reasons:

-Ionised species perturb the ion gauge measurement.

-The ion gauge does not discriminate the active species from the molecular and inactive N2 present in large quantity in the chamber (when the N plasma source operates, the background pressure in the growth chamber is in the range of 10 -6 -10 -5 Torr).

Thus in order to evaluate the flux of active N species, it is necessary to use growth rate measurement of planar GaN using RHEED. monolayer, resulting in a periodic variation of the intensity of the diffracted beam. Thus when the layer has a maximum flatness, the intensity of the spots is at the maximum and when the coverage is close to half of the monolayers, the intensity of the spot is at its minimum. Thus each RHEED intensity oscillation corresponds to the growth of one complete monolayer. By recording this intensity oscillation, it is possible to determine the average growth speed in-situ: V = n /t monolayers.sec -1 where n is the number of cycles of oscillations (number of monolayers) and t is the time corresponding to the particular number of oscillations.

NITROGEN FLUX CALIBRATION: The above mentioned RHEED technique can be used to calibrate the GaN growth rate and hence to deduce the corresponding active N flux when growth is limited by N. We used a GaN "template" substrate to grow GaN epilayers at 730 o C. By setting a random fixed amount of nitrogen flow, (by fixing the power and the SCCM), and by varying the Ga flux, it is possible to determine the stoichiometric condition (V/III ratio equal to 1) based on the following three scenarios:

1) When N / Ga ratio is greater than one, growth rate is limited by the Ga flux. Hence the growth rate increases by increasing the Ga flux. One thing to be noticed at this growth condition is that with increasing number of monolayers of GaN, the surface tends to get rougher. This is visible through the damping RHEED oscillations. The surface can be smoothened again by Ga and N fluxes alternately (by keeping only one of the shutters open), resulting again in sharper RHEED oscillations. This regime correspond to N rich growth conditions.

2) When N / Ga ratio is less than one, the growth is limited by the N active species. As the N flux is fixed, any change in Ga flux will not change the growth rate. Instead excess Ga is accumulating on the surface in the form of a metallic layer floating on top of the grown layers. As a result, no RHEED oscillations are observed. This regime is termed Ga rich (or metal rich) growth.

3) A N / Ga ratio equal to one corresponds to the stoichiometric condition. The amount of Ga and active N fluxes reaching the surface are equal. This is found by noticing that the growth rate saturates for any increase of Ga or N fluxes. At this point, it can be deduced that the quantity of the active N species is equal to the flux value measured for Ga using ion gauge for this cell temperature.

In this thesis, a N flux of 0.7 ML.s -1 (0.53 SCCM, 350 W, chamber pressure = 8.3x10 -6 Torr) and a Ga flux of 0.6 ML.s -1 are used to produce a V / III ratio of 1.1. Al flux is 0.15 ML.s -1 (8x10 -8 Torr).

Aluminium Nitride pre-deposition

GaN growth on bare Si (111) is known to produce SiN layer before growth of GaN. This is mainly because nitrogen bonds with Si faster/easier than forming GaN. It has been shown by TEM observations that SiNx is present between the NW and the Si substrate. This can prevent the epitaxial growth of GaN. As a result, NWs are tilted and have a random in-plane orientation. A common approach to overcome SiNx formation is to use Al-Si droplet gives rise to few AlN nanocrystals of platelet morphology 12 . The nanocrystals are N polar due to the strong tendency to fix the abundant Al atoms of the liquid droplet to the Si (111) dangling bonds. The crystals typically have a height between 10 -20 nm and a polygonal edge with a flat top 12 .

The surface thus obtained has a peculiar morphology [fig. 7]. It consists of N-polar AlN nanocrystals in epitaxy with the silicon substrate surrounded by a thin AlN layer of metal polarity.

GaN NW growth

GaN NW growth on AlN layers is studied in detail by many research groups 13,14,15,12,16 . Based on these results, there are four phases involved in growth of GaN NWs:

1. Incubation phase: Upon exposure of Ga and N flux (opening of shutters), GaN NW growth does not start immediately. There is a delay between the time when the fluxes are supplied and the onset of NW formation revealed by RHEED spots appearing on the screen. This duration is termed as the incubation time and could be represented as follows:

Tinc = TRD -TOS
where TRD is the time at which RHEED spots are observed and TOS is the time at which the shutters to supply the fluxes are opened.

2. Nucleation phase: Nucleation phase is a stage in which NW growth starts. The exact mechanism behind the formation of GaN NW nucleus is still under debate.

According to Consonni et al. 15 GaN first forms spherical islands which acts as a precursor for NWs. These spherical islands undergo a shape transition to for pyramidal structure which acts as the stable nucleus for NW growth.

3. Beginning of elongation phase: This phase mainly consists of elongation of GaN NWs. Growth is mainly supplied by the impinging active species on the NW sidewalls and the substrate surface. The adatoms diffuse along the sidewalls, reaching the top of the nanowires (favourable site for incorporation of adatoms) and participate in the axial growth of NWs, however it is true only if the diffusion length of adatoms is greater than the height of NWs 17,14,18 . In this growth stage of NWs, the elongation rate increases with time.

4. Axial and radial growth of NWs: Once the height of the NWs is greater than the diffusion length of adatoms, the axial growth rate stabilizes. At the NW growth temperature, axial growth is favoured at the expense of radial growth. However, all the adatoms cannot reach the top of NWs due to the increased height of NWs and there is a radial and axial growth happening simultaneously at this stage. Yet indeed, the incorporation rate of Ga adatoms is much higher on the GaN {0001 ̅ } facets (i.e. the top of NWs) than on the {1010} facets (i.e. on the NW sidewalls). The NW thus maintains the high aspect ratio.

EXPERIMENTAL CONDITIONS

After the deposition of the AlN layer, the GaN nanowires are grown in the temperature range between 760-810 o C in nitrogen-rich growth conditions. According to the V / III ratios and the growth temperature, the morphology of the NWs varies. These conditions may differ between each MBE setup, depending on the geometry of the growth chamber. The NW-2D growth regimes can be denoted in terms of Ga flux vs growth temperature. Growth diagram corresponding to our MBE chamber for a N flux of 0.7 ML.s -1 is shown in fig. 9.

The density and diameter of the NWs can be adjusted by varying the growth temperature and the V/III ratio. Of course, the length of the nanowires can be adjusted by varying the growth time.

There are different models which tentatively explain that the NWs grow on AlN pedestals 12,19 or in the edge-interface between AlN pedestal and the SiNx layer formed during the buffer growth 20,21 . However, the former idea is universally accepted and these different studies agree on the fact that GaN NWs grown after such a AlN pre-deposition step, are of very high verticality and have the in-plane orientation of the AlN pedestals epitaxially grown on the Si substrate 22,23,12 .

Most of our standard growth of GaN NWs are done at a V/III ratio of 1.1, a substrate temperature of 800 o C and a growth rate corresponding to 0.62 MLs.s -1 . The geometry of the chamber and the angle of incidence of the fluxes on the substrate can play a role on the morphology of the NWs 24,18 . To minimize their influence, the sample is kept rotating during growth (~10 rpm).

Characterisation of GaN NWs

Morphology of the NWs

We predominantly grew GaN NWs on Si(111) with AlN pre-deposition. It can be seen from the following SEM images that the growth with AlN produces highly vertical NWs. With AlN pre-deposition, NW are often found with a thin crystalline layer between the NWs. Top view-SEM images of the isolated NWs show their hexagonal cross-section. Cross-view shows highly vertical NWs with the presence of 2D layer between the NWs. These conclusions are in agreement with a previous study 23 . between the basal plane (0001) and the pyramid facets is about 58° which corresponds to the angle between (0001) and {112 ̅ 2 ̅ } planes as can be seen from the table in Appendix.

We used KOH solutions of concentration 0.4 mol.l -1 to etch the NWs for 10 minutes at 40 °C. All the NWs are found to be dominant N-polar. NWs before the KOH etch and after the KOH are shown in fig. 13. As it can be seen all the NWs have flat tops before the etching and sharp / sloped tops after the etching, which is a typical signature of N-polar NWs. to free A excitons (FXA) 28,29 and has a FWHM of 12.7

Photoluminescence measurements

meV. The 3.472 eV peak is relatively broad as compared to the best FWHM values found in literature (2.5 meV 6 , 0.5 meV 30 ). This can be possibly associated to a higher residual doping in our chamber. The 3.42 eV weak subbandgap emission may be present in the spectrum (however, its intensity is very weak). It may originate from excitons bound to planar defects 6 . The 3.45 eV sub-bandgap contribution is rather pronounced in our samples.

It is often reported in GaN NWs grown on Si (111). It has long been associated to the two-electron-satellite effect but the recent in depth analyses from T. Auzelle et al. 6 and Pfüller et al. 31 have unambiguously associated it to the presence of inversion domain boundaries (IDBs) in GaN NWs grown on Si substrates by MBE. Though it has been observed in GaN thin films as well, due to the random occurrence of the defect, it is tough to spot. In the case of NWs, as they are thin, the IDBs can be readily spotted using STEM observations. Auzelle et al. 6 have observed such IDBs in [112 ̅ 0] and [11 ̅ 00] zone axis, [112 ̅ 0] being more favourable due to the lattice spacing between Ga and N atoms 6 . NWs removed from the previous sample, grown on Si (111), were observed by STEM. As suggested by the presence of the 3.45 eV peak in their PL spectrum shown in fig. 14, several NWs contain a polarity inversion domain. The inverted domains (of Ga polarity) could not be identified with the KOH etching method due to the fact that when present, these Ga-polar domains represent a minor part of the NW volume. The inverted domains seem to be located in the core of the NWs and they extend from the bottom to the top. Therefore, the description of NW nucleation presented in section 2.2.4 is incomplete. It seems possible that at the early growth stage, Ga-polar NWs nucleate besides the N-polar NWs. A discussion about this will be presented in the photoluminescence section of chapter 4, when we compare the NWs grown on Si (111) with those grown on silica substrates.

INVERSION DOMAIN BOUNDARIES

Summary

In this chapter, we have explained in detail the growth procedure that we use to obtain GaN NWs on Si (111) substrates with AlN pre-deposition and we have presented the state-of-the art results with this approach.

The experimental steps involved in self-induced growth of GaN NWs on Si (111) are discussed. Surface preparation, pyrometer and flux calibration using RHEED are explained and similar procedures will be applied in the following chapters. We have also presented the different characteristics of NWs grown on Si (111) in our MBE chamber. Their surface morphology, in-plane orientation, verticality, polarity and PL measurements are reported. We have shown that IDBs are commonly found in these NWs with a signature at 3.45 eV in the PL measurements. These samples and their characteristics will be used as references for comparison with NWs grown on other novel substrates. AIC-Si as a substrate for gallium nitride nanowire growth

Introduction and objective

In semiconductor materials industry, bulk crystalline substrates are used for epitaxial growth of materials where these substrates act as a support, as well as an epitaxial seed layer for growth of high quality active layer. Usage of such bulk substrates results in material loss as the main purpose lies in the crystalline structure of the surface of the substrate. On the other hand, polycrystalline thin Si films can be produced on large scales at low cost, low processing temperatures (compared to the monocrystalline Si bulk substrates) and using simple deposition processes (compared to 2D materials). In this project we would like to fabricate thin (in the order of tens of nm), crystalline Si films with wide grain size (in the order of µm) having a good <111> fibertexture, i.e., all the grains having a common <111> direction perpendicular to the substrate (will be explained in detail at section 2.1.3.1). In our work, reported later, these pc-Si layers act as a template for growth of GaN NWs by MBE.

I prepared the AIC-Si samples in Unite Mixte Saint Gobain Recherche -CNRS, Aubervilliers.

Means of obtaining thin polycrystalline-Si layers

As seen in the previous section, Si is one of the main hetero-epitaxial substrates being actively explored for growth of GaN NWs. GaN-on-Si has been emerging so fast in the past decade that it could be considered as a matured technology. In most of the cases, bulk crystalline Si wafers with a (111) orientation are used as substrates by overcoming the lattice mismatch and the thermal expansion coefficient mismatch by using intellectual property (IP)-protected buffer layers 1 . This is less interesting when we take into account the large material loss (because of the bulk substrate) and the inaptness to the future scope of producing flexible devices.

It is therefore important to use a means to obtain thin Si layer which could act as a thin-flexible-supports for high quality material growth.

One approach is to grow Si thin film epitaxially on a crystalline Si wafer with a sacrificial layer and detach them to produce free standing thin substrates 2 . Though it enables to produce large area monocrystalline Si templates, the process is limited by the size of the Si wafers (not individually scalable) and introduces significant cost and complex processing steps 3 . Another common way to obtain thin Si layer is by means of plasma enhanced chemical vapour deposition (PECVD). However, this method produces amorphous Si layers or polycrystalline Si layers with the grain size in the orders of nm 4 . Though this could be sufficient for NW based devices due to the small footprint of NWs, it is impossible to scale to micro-structures (micro-rods or micro-disks having footprint in the orders of µm) as it might induce defects from the grain boundaries of Si. With CVD at high temperatures (>700 o C), it is possible to grow polycrystalline Si directly on any foreign substrate 4 . The size of the grains is controlled in the nucleation phase and are typically in the range of 2-20 µm. The layers obtained often present a [110] fiber texture 5 which is due to the fact that [110] oriented grains grow faster than others.

Thus the texture quality is much stronger as the thickness of the layer increases. However, it is not an indication of the layer quality as twinning defects are commonly observed in such layers with high fiber texture quality 4 .

Another important method to obtain crystalline Si films is by deposition of amorphous Si (a-Si), and a subsequent annealing step to crystallise Si. This method involving a simple annealing step of a-Si for many hours, is referred by the term "Solid Phase Crystallisation" (SPC) 6,7 . Deposition of a-Si can be obtained in many ways like CVD 7 and sputtering 8 . The layer is then annealed for several tens of hours at a temperature between 550 o C and 700 o C 9 . By this method the higher the temperature, the faster the crystallisation, but also smaller the grain size 6 . This method of crystallisation is generally used for layers having a thickness of 1 µm and the grains have an aspect ratio close to unity and feature a [111] fiber texture. But the main drawback with this method is the long annealing duration to get good quality fiber-texture layer, calling for a compromise between layer quality and the duration of the process 6,9 .

Alternative to SPC, in the photovoltaic research industry, exists a well-known process to produce high quality polycrystalline thin-Si films (active layer) on top of another c-Si thin films (seed-layers) 10 . This is done in order to separate the doping issue from the morphological quality of Si layers. These seed layers are not of perfect optical quality for photovoltaics but they consist of grains of large lateral dimensions 4 . Among the methods to produce these seed layers, we can first mention about the method called "Recrystallization", where the a-Si is in contact with a c-Si layer and annealing is performed in a tube furnace where radiation is focused in a line on the surface of the substrate 4 or by a laser 11 . These two means provide layers with a [100] or [111] texture depending on the initial orientation of the initial crystalline domains and the thickness of a-Si films. These methods produce grains of very large lateral extension (from few µm to mm) for very thin layers (25 nm). They are much faster than the SPC, however require expensive equipment to conduct crystallisation 12 . The second method is called metal induced crystallisation (MIC) where annealing an a-Si in contact with a metal layer greatly reduces the thermal budget, up to a factor of 66 (in contact with Al) 9 .

Aluminium induced crystallization of Si (AIC-Si)

In 1970s it was discovered that the presence of a metal in contact with an a-Si layer greatly accelerates the kinetics of the SPC 21 . Many metals were subsequently tested with success 22 involving Al is called either 'Aluminium Induced Crystallization of Si (AIC-Si)' or 'Aluminium-Induced Layer Exchange (ALILE)'. This process is also particularly interesting for photovoltaic industry as the bulk crystalline Si wafers are indicated to take the large part (~50%) of expense 23 to each other, is termed fiber-texture. Fig. 2 shows a graphical representation of a <111> fiber textured layer.

In our case, the target is to have domains of lateral dimensions in the order of µm, much higher than the footprint of the NWs so that each NW will grow on a locally-monocrystalline surface. This will enable to have monocrystalline AIC-Si islands when moving to patterned AIC-Si layers (will be explained in detail in section 3.4.2).

Proof of concept for GaAs NWs

A preceding PhD student, Y. Cohin, identified that this process could be of significant interest to NW growth community 24 . The thin polycrystalline Si layer obtained through AIC-Si has a <111> fiber texture which as substrates for epitaxy of NWs.

Growth of Ge nanowires on AIC-Ge substrates

Along with silicon, germanium is one another material which is vastly explored for synthesis using AIC process 26,27 . Toko et al. 25 , have shown a direct fabrication process (without a transfer) of vertical Ge NWs on a flexible plastic support by the use of AIC-Ge as substrate. The Ge substrate with a [111] fiber texture was fabricated at a temperature of 325 °C. Ge NWs were obtained by CVD at 360 °C for a growth time of 10 minutes using vapor-liquid-solid (VLS) growth mechanism. Au nanocolloid particles are used as the catalyst for VLS growth. For growth on substrates like a-Ge, polycrystalline Ge (without fiber texture), SiO2 and directly on plastic, the NWs grew with random alignment. However, on [111] fiber textured polycrystalline Ge, the NWs grew with a very high verticality along the [111] crystal direction, indicating that the fiber-texture quality of the AIC-Ge substrate has a strong influence on the orientation of the NWs. TEM analysis of the Ge NWs obtained on

Objective of this work

With this work, we would like to show that AIC-Si could act as a template for GaN NW growth on an amorphous substrate. In order to improve the verticality of the NWs, it might be required to reduce the surface roughness and enhance fiber-texture quality of the AIC-Si layer. It will be appropriate to explore the standard growth conditions used with the conventional Si(111), substrate for AIC-Si layer. Finally, it will be appropriate to obtain selective area growth of NWs on AIC-Si with respect to amorphous material which can be the support of the AIC-Si layer. For this purpose, a silica layer will be used as a support. AIC-Si patches will also be fabricated and used to obtain vertical NWs on defined spots of the substrate.

Synthesis of AIC-Si

Thermodynamics of AIC-Si formation

In order to optimise the layer quality of the final polycrystalline silicon (pc-Si), it is necessary to understand in detail the mechanism of AIC-Si formation. In this section we will see the detailed understanding of the AIC-Si formation found in literature. These studies form the basis from where we started our research in optimising the AIC-Si layer.

The more detailed understanding of the AIC mechanism involves the study of a-Si/Al interface. By thermodynamic considerations, amorphous materials strive to crystallise because the Gibbs free energy of crystalline state is lower than the amorphous state 9 . But due to the strong covalent bonding (3.30 eV for Si-Si) present between the atoms, exists a high activation energy for the conversion of a-Si to c-Si. Usually the overcoming of this barrier is achieved by thermal activation and requires a temperature of about 700 o C for Si.

When an amorphous semiconductor is in contact with a metal, this thermal activation temperature decreases indicating the activation barrier is reduced somehow. Experiments by Hiraki et al. showed that the strong covalent bond in semiconductors can be weakened at the interface by the presence of a metallic layer 30,31,9 . The a-Si atoms present at the Al/a-Si interface undergo covalent bond weakening by Coulomb screening effect, and this promotes an enhanced mobility of the Si atoms 32,9 . However, it has been proven that nucleation of c-Si is not possible at the Al/a-Si interface at very low temperatures (<130 o C) 33,9 . The Al layer deposited by sputtering, is usually polycrystalline with randomly oriented grains 34,33 . Nucleation of c-Si happens at the grain boundaries of Al layer. Through TEM experiments it has been shown that wetting of high-angle Al grain boundaries by a-Si, happens at temperatures slightly below 140 o C and the nucleation of c-Si starts at 150 o C at high-angle Al grain boundaries(GBs) 33,9 . This migration of a-Si Al/a-Si ratio is less than 1, there is a continuous layer of pc-Si layer on which new pc-si islands form. This is observed as a two-level pc-Si layer or pc-Si continuous layer and island growth on top 35 . The thickness ratio of amorphous layer to the metal layer, in general plays an important role in the getting either a continuous layer or islands of crystalline semiconductor layer in both the case of Si and Ge layer 35,36,37 .

Annealing process

As seen in the previous section, the crystallisation of AIC-Si consists of two phases: 1. Nucleus formation and 2. Incorporation of a-Si atoms into this nucleus thereby extending the grains laterally until all the a-Si is consumed or the grains coalesce 34 . Each of this phase happens only on annealing the layers. As a result, annealing temperature and time are important parameters in the AIC-Si formation process 9 . Higher annealing temperature (~500 °C) produces large density of c-Si nuclei rapidly which understandably reduces the grain size due to coalescence of grains 38 . Lower annealing temperature needs longer time to crystallise 100% (amount of material) of the a-Si and indeed the throughput is reduced 38 . Hence a compromise is needed to have maximum grain size with lesser time possible to crystallise an entire a-Si layer. For an AIC stack with 150 nm a-Si and 50 nm Al, it has been observed by x-ray diffraction that the integrated intensity of Si (111) peak is highest after 5 hours of annealing at 250 °C and does not change even after annealing for 25 hours 9 . It is an indication that the a-Si reaches close to 100% crystallisation and almost no new nucleus was formed after 5 hours of annealing. We have followed the same thermal budget (250 °C for 5 hours) as a first step to create and extend the c-Si nuclei considerably. And further to accelerate the later extension of grains and achieve a 100% crystallinity of the Si layer, we anneal the layers at 500 °C for 30 minutes as a second step. The maximum temperature is chosen as 500 °C as it has been shown that solubility of Si in Al is negligible at low temperature (~200 °C) and increases considerably for temperatures above 500 °C. For temperatures higher than this and close to the eutectic temperature, Si atoms diffuse into the Al layer itself rather than at the Al grains and form pc-Si inside the Al grains 9 . Hence the maximum crystallisation temperature is limited to 500 °C in our case. We call this annealing process as '2-step annealing of AIC-Si'. Moreover, the same technique has been shown for SPC with the same motivation; A two-temperature annealing was followed to greatly reduce the total annealing time (compared to low temperature annealing) and to increase grain size (compared to the higher temperature annealing / RTA) 6 .

Essence of intermediate oxide layer

Intermediate oxide layer plays a significant role in the fiber texture quality of pc-Si layers 39,40 . Kurosawa et al. 39 , have reported that lack of an intermediate oxide (IMO) layer results in randomly oriented pc-Si layer whereas with sufficient thickness of intermediate oxide, Si(111) fiber texture was produced. The intermediate oxide layer is explained to act as an inter-diffusion barrier for a-Si at low annealing temperatures. Depending on the thickness of silica, the group demonstrated Si(100) and Si(111) fiber textured pc-Si layers. In their experiments, the oxide layer was prepared by exposing the sample to air for different duration of time. Thin layer of IMO (5 minutes exposure to air) at the interface of a thick a-Si/Al (100 nm) stack produces Si(100) oriented layer while thicker layer of IMO (24 hours exposure to air) for the same stack produces Si(111) oriented pc-Si. Further results reported by Okada et al. 40 showed that there is a presence of an optimal thickness IMO layer for specific thicknesses of AI-Si stack above which the fiber-texture is again affected. In their case, for a a-Si/Al stack of 100 nm each, 10 nm is the optimal thickness of IMO to produce a factor close-to-unity of 111 fibertextured c-Si layers.

Apart from controlling the orientation of the fiber-texture, IMO thickness also plays a role in controlling the growth rate of c-Si 40 . It shows that IMO acts a permeable layer in diffusion of a-Si into the Al grain boundaries.

In fact, it has been shown that IMO plays a crucial role in controlling the dissolution of a-Si in Al at the a-Si/Al interface at high annealing temperatures (close to the eutectic temperature) 9 .

Preparation and optimisation of thin Si layers

Experimental section

In this section we will explain in detail about the layer preparation process, right from deposition to crystallisation of Si and removal Al to have the final pc-Si layer ready for growth NWs on them. In our case we have used a growth support consisting of 25 nm SiO2 layer thermally grown on Si (100) quarters to replicate the properties of a glass or amorphous support. hours and the sample is transferred to Rapid Thermal Annealing (RTA) oven for the second level annealing at 500 °C for 10 minutes. As there is no significant difference in the key outcomes (grain size, surface roughness) between samples prepared using these two different methods, for most of the samples two step annealing is performed in the normal oven itself.

As seen before this annealing steps results in pc-Si with the exchange of layers. As a result, Al is found on the surface with the pc-Si moving under on to the amorphous substrate.

SURFACE TREATMENT

The resulting Al/pc-Si stack is dipped in 5% HF for approximately 10 seconds to remove the Al layer. We have seen that with more dipping time, there is an under etch resulting in removal of amorphous layer under the pc-Si.

Optimisation of the layer

Possible minimum thickness for continuous c-Si layer

In order to use AIC-Si layers as substrates for NW growth, one of the requirements is to have thinnest layer possible. Thin layer could provide easy transfer after NW growth and transparency of the Si film if necessary for the device operation. For these reasons, to synthetize thinnest a-Si layer with Al thickness of 10 nm, we used a-Si thicknesses of 5 nm, 10 nm and 15 nm. Samples were prepared using the standard sputtering conditions and standard annealing conditions (250 o C for 5 hours, 500 o C for 30 minutes).

From SEM observations it can be seen that for a thickness of 10 nm Al, 15 nm a-Si is required to have a good final pc-Si coverage. By imageJ analysis, for this sample the area coverage of pc-Si layer is estimated to be close to 82 %. Hence we concluded that a thickness of 15 nm of a-Si can be used for AIC-Si for our NW growth experiments. However, this thickness of a-Si presents the problem of 2 level-pc Si formation. There is first layer

GaN nanowires on AIC-Si

Growth of GaN nanowires on continuous AIC-Si layers

In order to test growth of GaN NWs on AIC-Si template, we took a continuous layer of pc-Si fabricated with the standard IMO and with a-Si and Al stack thicknesses of 15 nm each on top of 25 nm-SiOx/Si(100) support.

On annealing using standard conditions (250 °C for 5 hours and 500 °C for 30 minutes), recrystallization is completed and the Al layer was etched away by dipping in 5% HF for 10 seconds. This pc-Si template was introduced into the MBE chamber immediately after the HF etching. After the standard in-situ preparation of substrate in the introduction chamber (Baking for 10h to eliminate adsorbed gases) and in the growth chamber (850 °C for 10 minutes to remove the surface oxide if any), growth of GaN NWs was conducted using our standard growth conditions for GaN NWs which include AlN pre-deposition. Temperature of the pc-Si sample is read by pyrometer that is calibrated using a Si(111) wafer by the appearance of 7 by 7 reconstruction at 830 °C.

A schematic of our standard growth procedure for GaN NWs is shown in fig. 18. This procedure has been optimised for growth on Si(111) bulk monocrystalline wafers. Al is deposited at 620 o C for 20 seconds and it is exposed to N for 1 min at the same temperature. After the substrate temperature is ramped up, GaN NWs were grown at 800 o C for 3 hours with a V/III ratio of 1.1. This procedure is traditionally followed for growth of With respect to the AlN step, there could be two different factors involved: firstly, it can be simply due to the extra roughness introduced on pc-Si surface due to the AlN pedestals. This is not very likely considering the tilt of the NWs and the size of these AlN pedestals.

Secondly, it could be due to the possibility of Al droplets dissolving the thin pc-Si layer locally. Indeed, the substrate temperature for Al deposition is 620 o C, i.e., well above the eutectic temperature of 577 o C for Al-Si.

On supply of N2, this could lead to formation of AlN in random direction since the underlying support is the amorphous SiO2. A similar phenomenon is reported for growth of Si NWs on thick AIC-Si template (50 nm) with Au droplet 28 .

Characterisation of the NWs

Verticality of NWs

In order to have a quantitative analysis of the deviation from verticality of NWs, we used pole-figure measurements by XRD. GaN (101 ̅ 1) planes which are at 62° from the (0001) planes are used to detect the in- The position of this ring (71°) shows that majority of the Si grains have (111) planes oriented normal to the substrate (in agreement with the grazing incidence analysis. Cut-line profile taken by integration of intensity the composition of droplet is changed 28 .

In order to verify this, we used the Si layers fabricated with 5nm a-Si / 10nm Al stack, where the AIC-Si is an incomplete layer (present in the form of islands). To observe the substrate surface below the NWs, we grew shorter and less dense GaN NWs using relatively a higher temperature (805 °C). To understand that the difference is only due to the AlN pre-deposition step, we fabricated samples with and without the AlN predeposition step.

As it can be seen in the SEM images, the AIC-Si layer with the AlN pre-deposition step consumes the Si thin layer and forms pedestals on which the NWs grow vertically. Based on the facets seen on top-view SEM images, we also observed that the NWs on same pedestals are in-plane oriented (fig. 22 (d)).

Polarity of NWs

We used KOH etching method to determine the polarity of the NWs. This method is commonly used in determining the polarity of 2D layers 45 . KOH solution etches selectively the N-polar end of GaN layer and is inert

Surface deformation

In both the growth procedures, after the growth of GaN NWs on pc-Si, it can be observed that there are some stone-like structures on the substrate only on top of the pc-Si layers. In most of the cases, no NWs were observed on this stone-like structures.

This undesirable evolution of the AIC-Si layer might be due to the high growth temperature. To verify this, the AIC-Si layer was heated in MBE chamber without the growth fluxes (shutters closed and cells at standby temperature). The layer was annealed in the MBE chamber with the substrate temperature at 800 o C for 3 hours.

Fig. 25 shows that, indeed, the AIC layer morphology is strongly affected by this thermal treatment. Through EDS analysis we found that the stone-like structures are blocks of Si. No contamination or signal for any other elements are detected. This clearly shows that the AIC-Si layer deformation observed during the growth of GaN NWs is related to the high growth temperature (greater than 800 °C), and not to the interaction with Ga, Al or N fluxes.

At temperatures greater than 600 o C, normal crystallisation (without a necessity for metal interface) of a-Si can happen (Solid-phase-crystallisation) 47,9 . In order to understand the effect of this surface deformation on the fiber texture quality of the layer, we did the 2theta/omega comparison of the layer before and after thermal treatment at 800 °C. There is no change in the peak corresponding to (220) orientation, however the peak corresponding to the {111} planes is relatively more intense.

Though the exact reason why this deformation is happening is not clear, we are wondering whether it might be due to the eutectic alloy formation due to some tiny amount of Al present in the pc-Si layer. We recall that the initial stack for AIC process contains an Al layer. Another explanation could involve the recrystallization of the second level of AIC-Si observed in the 2-level MIC layers. It has been shown that during the formation of twolevel pc-Ge layer, the second layer is present without the fiber-texture 36,37 . For AIC-Si as well it has been shown

Growth on patterned AIC-Si

Many groups investigate selective area growth (SAG) for elaborating GaN nanostructures with a reduced defect density as compared to 2Dfilm. Additionally, SAG is important in elaborating controlled micro-or nano-structures at precise locations. On crystalline bulk substrates, a common method to obtain SAG is to use dielectric mask deposited on the substrate surface and patterned with small openings, called 'holes', to expose the crystalline substrate. This approach is not practical for pc-Si layer fabricated as this might expose grain boundaries of the Si layer at the patterned holes. Hence for a pc-Si layer, it is important to produce separated, individually monocrystalline patches of Si.

Preliminary results for selective growth

In order to have selective area growth (SAG), it is necessary to optimise the growth conditions of NWs, in terms of temperature, so as to avoid the growth on SiO2 support and obtain NWs only on AIC-Si. We used a simple method to prepare a SiO2 surface partly coated with the AIC stack by hindering the flux of sputtered materials for a part of the wafer by using a piece of another wafer, attached to the AIC-Si support (SiO2/Si(100)).

Once the sputtering is done, the masking wafer is removed and crystallisation and removal of Al by HF is done.

It resulted in AIC-Si on one half of the wafer and the other half of wafer having the SiO2 support, to test the growth selectivity on AIC-Si templates, without the need to go for expensive patterning for initial results.

For the growth of GaN NWs on these templates, a V/III ratio of 1.1, at a growth temperature of 800 °C, yielded NWs on both the AIC-Si layer and the SiO2 support. However, it can be seen that the NWs in AIC-Si layer have higher density and height as compared to the ones on SiO2. We tried growing at higher temperature and it can be seen that at a growth temperature of 810 °C, almost no NWs were present on SiO2 and non-coalesced, vertical GaN NWs were obtained on AIC-Si layer. Thus for the future growth on e-beam patterned AIC-Si, a growth temperature of 810 °C was used with a V/III ratio of 1.1.

Fabrication of AIC-Si nano-islands by lithography

To obtain the nano-patches for SAG of NWs, we followed the below steps:

1. Deposition of PMMA and MAA by spin coating on the amorphous support on which the pc-Si and NWs are need to be grown.

2. E-beam patterning of the PMMA layer and development of PMMA layer.

Vertical GaN NWs on AIC-Si nano-islands

We grew GaN NWs on these nano-substrates prepared by patterning of AIC-Si stack (mentioned in the previous section). After the Al removal by HF etching, the sample was introduced into the MBE chamber directly.

Usual thermal treatment at introduction chamber and main chamber was carried out. Then, NWs were grown at a temperature of 815 o C to get selective area growth. Temperature of the pyrometer was calibrated using the standard calibration technique with a Si(111).

We observe proof for vertical growth of GaN NWs on monocrystalline Si islands fabricated using AIC-Si process. The growth was highly selective with NWs present only on the c-Si islands and almost no growth on the silica surface present around the c-Si patch. For an island width of approximately 90 nm, we see only 1NW per island. However, the yield of such a growth is quite low and needs further optimisation of the growth conditions.

When the island size is larger than the footprint of the NW, we can see that more than one NW nucleate and grow on the islands. The NWs grown on the edges of the islands are mostly tilted. As opposed to the growth on continuous layer of AIC pc-Si (explained in section 2.3.3) the c-Si in an island form, is not deformed and are intact even after a growth of 3h at temperature greater than 810 o C. This could be because of a couple of reasons:

1. Because of their small size, the c-Si patches have necessitated longer annealing at low temperature (250 o C for 20 hours). This step could have crystallised 100% of the a-Si with less residual incorporation of Al inside the c-Si phase, as compared to the situation at high annealing temperature (500 o C). We recall that high temperature annealing (~500 o C) incorporates higher percentage of Al in pc-Si layer compared to annealing at 250 o C 9 . For AIC layers containing more Al inside, when we increase the temperature above the Al-Si eutectic temperature after the removal of Al layer by HF, there could be a significant dissolution of pc-Si forming a liquid Al-Si alloy. From the eutectic diagram for Al-Si it can be seen that at 800 o C, a given amount of Al can dissolve an amount of Si twice larger.

2. When the initial AIC stack is patterned into small patches, the formation of a second level of c-Si during the crystallisation process is probably avoided, because the grain being crystallised has free boundaries.

Moreover, we have used an Al layer of 6 nm in the AIC stack for the patterned samples. We have shown previously that with such a thin Al layer, the second level of AIC-Si (island + continuous layer) does not form. We would like to draw attention to the point that in the presence of 2-level AIC-Si, the top level does not adopt the fiber texture of the bottom level. We suspect that the AIC-Si layer could recrystallize at higher temperatures like 800 o C (growth temperature of GaN) thereby consuming the fiber-textured lower level of pc-Si as well. This calls for additional inspection of the layers fabricated with 6 nm Al/15 nm a-Si stack, which is not covered in this work yet.

Conclusion and future scope

In conclusion we have shown that epitaxial growth of GaN NWs by MBE is still possible on c-Si fabricated by AIC method. On continuous pc-Si layer, problem of deformation is observed though it is not completely clear on the reason behind it. However, on patterned c-Si islands, the layers were intact and we were able to obtain vertical GaN NWs. The selectivity of GaN NW growth on c-Si islands is high with respect to the silica mask layer.

We showed that for Si island size comparable to the footprint of NWs, 1NW per island is achievable. Though there is a long way in obtaining a good yield for 1NW growth per island, it will help in fabricating heterostructures based on GaN NWs on flexible supports, on a substrate (AIC-Si) similar to well-examined and understood bulk Si crystalline substrates. Back contact can be provided by use of 2D materials (like graphene 49 ) or metallic layers.

Indeed, it has been shown that the MIC layer can be grown in the same way either on glass or on a metallic surface [START_REF] Gall | Conf. Rec. Twenty-Ninth IEEE Photovolt. Spec. Conf[END_REF] . Thus there is a high scope for using AIC-Si layers as the main substrate for device fabrication with GaN NWs.

Introduction and objectives

Amorphous Substrates

Self-induced GaN NWs are grown traditionally on various crystalline substrates such as Si(111), Si(001) 1,2 , sapphire 1,3 , diamond 4 and SiC 5 . As previously mentionned in section 1.3.1 in order to contact efficiently an ensemble of NWs and fabricate devices, it is important to grow vertical NWs. This requirement is most easily met when monocrystalline substrates are used since wurtzite GaN NWs generally elongate along the [0001] axis which aligns respectively with the [0001] and [111] direction of the hexagonal and cubic substrates. This epitaxial alignment is not systematic: several studies 2,6 using monocrystalline Si substrates have reported the formation during growth of an amorphous layer due to the reaction of the substrate surface with the atomic N flux. This amorphous surface layer can induce some NWs to grow in random directions while the others grow vertically following an epitaxial relationship with the crystalline Si lattice. The orientation of the NWs that nucleate on the amorphous areas depends on the roughness of the amorphous layer. Thus by having a flat amorphous substrate, it might be possible to grow vertically aligned NWs.

From the applications viewpoint, there is a strong interest to develop a procedure for the wellcontrolled NW growth on non-crystalline substrates such as, for example, glass. Glass substrates can be produced in large areas and are widely available. The transparent nature of the substrate with almost 0% absorption is very appealing for opto-electronic device fabrication. Thus amorphous substrates could bring new flexibility or functionality as well as cheaper production costs, especially for devices such as LEDs and solar cells.

Objective

An important issue is to clarify whether a high yield of vertical NWs can be obtained on an amorphous surface where the epitaxial alignment of the same type on a crystalline substrate cannot take place. This is the objective of the present study. To this aim, we elaborate GaN NWs on silica and we compare them with NWs obtained on Si (111) substrates in terms of verticality, in-plane orientation, polarity and photoluminescence properties. 4.3 GaN nanowires on silica

Thermal and fused silica substrates preparation

In this chapter we used three types of substrates for NW growth study on silica: monocrystalline Si (111) wafers as reference sample, thermal silica obtained by dry oxidation of Si (100) wafers and bulk fused silica wafers. The thermal silica was fabricated by thermal annealing of Si (100) substrates under a N2 -O2 gas mixture to obtain a 25 nm thick layer, as measured by ellipsometry. This thickness should be sufficient to prevent any possible epitaxial relationship between the Si (100) wafer and the GaN NWs grown on the silica surface. Before introduction into the oxidation oven, the Si (100) substrates were chemically cleaned using solvents and a short diluted HF 5% bath to remove any residues and native oxide as described in section 2.2.1.1. Same cleaning procedure was performed for Si (111) substrates before introducing in MBE for growth. Bulk fused silica wafers were used for the NW growth to exclude all influence of an underlying crystalline order. To assist the pyrometry readings, a 100 nm-thick layer of Ti was deposited on the backside of the fused silica wafer. No additional cleaning was carried out on thermal silica and fused silica wafer before introduction into the MBE. Finally, before growth, all the samples are in-situ outgassed overnight in the introduction chamber, followed by a hightemperature annealing at the growth chamber (detailed in section 2.2.1.1).

Growth of GaN NWs

In this study, the NWs were grown by plasma-assisted molecular beam epitaxy (PAMBE) under standard growth conditions elaborated in section 2.2.1.3. During NW growth, the substrate temperature was measured using an infra-red pyrometer and maintained constant at 800 °C unless otherwise specified. The absolute temperature was calibrated by observing the transition from a 1x1 to a 7x7 reconstruction of a Si (111) surface during sample cooling. To have consistent pyrometer measurement when the surface starts to be coat by GaN, we did not apply any correction of surface emissivity for the different substrates. When not specified, the GaN NW growth time was 3 hours. A few samples were elaborated using shorter growth times in order to reduce or suppress the shadowing or coalescence effects.

For comparison of NW quality, growth on Si (111) is performed with AlN pedestals formed prior to GaN NW growth, as described in section 2.2.1.3. We would like to recall that this procedure prevents the formation of a SiNx layer between the GaN NWs and the crystalline Si substrate and thus warrants that all the GaN NWs are grown in epitaxial relationship with the Si substrate.

Fig. 4 shows the as-grown NWs observed by scanning electron microscopy (SEM). The NWs grown on the amorphous silica (thermal silica in Fig. 4b, fused silica in Fig. 4 (c) show a remarkable vertical orientation.

This verticality is comparable to that obtained on the crystalline Si (111) surface (Fig. 4 (a)), although no crystal lattice underneath can guide the NW growth on silica samples. The top views in fig. 4 even reveal that less parasitic structures or strongly tilted NWs are present on the silica samples as compared to the reference sample. We also observe that the NWs on the amorphous surfaces are significantly shorter than on crystalline Si (Fig. 4 (d)-(f)). The diameter and width of NWs on Si (111) and thermal silica are 45 nm, 760 nm and 30 nm, 450 nm respectively. This is due to the high incubation time (40 minutes -measured by following the evolution of the RHEED diffraction patter) on thermal silica as compared to on Si (111) (5 minutes) and will be discussed later. Due to the high resistivity of the fused silica wafer, the direct SEM observation was difficult on this substrate because of large charging effects. To solve this issue, a 3-5 nm thin layer of Pd/Au was conformally deposited on the NWs by plasma deposition before observing these samples; this does not affect the NW morphology, but the SEM image (Fig. 4 (f)) is slightly blurred.

We wanted to study the nucleation and the initial stages of the NW growth. As mentioned before, the onset of the RHEED diffraction spots, characteristic of NW formation, was observed after about 5 minutes and 40 minutes for growth on Si (111) and thermal silica respectively. These delays correspond to the so-called incubation time which precedes the first nucleation events. The NWs grown on amorphous silica are shorter than those on crystalline Si (111) after growth of same duration (see Fig. 4). This is partially due to the longer incubation time on silica. However, another major difference between these samples is the NW density. The density on Si (111) is lower than on thermal silica (Fig. 4 (a)-(b)). The local environment of a given NW has a strong impact on its elongation rate, since the species supplied are shared between the NWs. The higher densities obtained on the silica substrates could also explain the shorter NW lengths.

Growth evolution

To investigate NW nucleation in more details, we observed GaN NWs grown on thermal silica for different times. authors observed dense arrays of NWs on the amorphous surface, while a rough compact layer was formed on sapphire, i.e. the NW formation was more favourable on the amorphous layer. This result contrasts with our study where the crystalline order at the Si surface seems to help NW nucleation.

In our case, the silica surface does not seem to present defects activating nucleation and the crystalline order at the Si surface appears more favourable for nucleation. This observation agrees well with several other studies where SiOx is used as a mask on crystalline Si, enabling selective growth on the exposed Si areas 12,13 . The absence of any deposit for exposures shorter than 40 minutes indicates that during the incubation, the incident Ga and N species mostly desorb from the silica, except a small fraction which accumulates as mobile adatoms.

The question whether the density of adatoms saturates very quickly (i.e. in a time much shorter than the This linear trend is not observed (Fig. 5). We This lowers the adatom density and strongly reduces the probability of further nucleation events, which explains why the NW density tends to saturate at longer growth time. However, the fact that new NWs continue to nucleate at time delays exceeding the incubation time, certainly contributes to the NW length dispersion which is observed in fig. 4. A third hypothesis which would not necessitate this slow increase of the adatom density, could involve chemical or morphological modifications of the silica surface exposed to the fluxes. In particular, the species produced by the N plasma may be responsible for such modifications which could create favourable sites for GaN NW nucleation. The incubation time would be the critical time necessary to form these nucleation sites. To elucidate this interesting question (what really happens during the incubation), more investigations would be necessary but are out of the scope of the present work. In the following, we do not evidence these possible modifications at the substrate surface. To quantitatively compare the verticality and the in-plane orientation of the NWs grown on fused silica, thermal silica and Si (111) surfaces, X-ray diffraction pole figure measurements were performed (Fig. 6a). They were recorded using the Bragg reflection of GaN (101 ̅ 1) planes. All the pole figures present a high diffracted intensity at the sample inclination ψ= 62°, indicating that the c-axis, i.e. the growth direction of the wire, is normal to the surface of the silicon or silica substrates, consistently with the previous SEM observations (Fig. 4).

As evidenced by the uniform ring in the pole figures of NWs grown on thermal silica and fused silica (Fig. 6a), on fused silica, and of 0.5° on Si (111), the latter being limited by the experimental resolution. Although the crystalline Si substrate leads to the best NW verticality, the angular distributions on the silica substrates are surprisingly narrow. The remarkable NW verticality obtained on the amorphous substrates contrasts with previous reports 14 which show that non-vertical NWs can form at the early growth stage and is consistent with the report by Stoica et al. 6 . Before NWs can grow, GaN crystalline nuclei must form on the silica surface. It is very likely that these GaN nuclei can minimize their energy by developing a (0001) facet at the interface with the silica surface. We confirmed this hypothesis by observing a cross-section of the interface between fused silica and NWs by TEM (explained in the following section). We have observed the NWs by transmission electron microscopy (TEM) using a 200 keV FEI Titan Themis equipped with a probe aberration corrector and a super-X Bruker silicon drift detector (effective angle of detection of 0.98 steradian). Cross-sectional samples have been obtained using a Dual beam FEI Scios Focused Ion Beam (FIB) system, after carbon deposition.

NW base has wurtzite structure with no stacking faults or defects and the (0001) facet rests directly on the oxide layer (Fig. 7 (b)). Before preparing the lamella by FIB, the sample was coated with carbon and locally, platinum. The C layer is about 50 nm high and appears with a bright contrast in fig. 7 (a) while the Pt layer is thicker and appears darker. The substrate surface around the NW, magnified in fig. 7 (c), presents a dark contrast. The composition of this two-dimensional (2D) surface layer is identified as GaN by energy dispersive Xray spectroscopy (EDX) analyses (Fig. 7 (e)-(f)). The presence of such a layer surrounding the NWs on top of an amorphous oxide was already mentioned by Stoica and co-workers 6 , who stated that this GaN layer was also amorphous. However, with an adequate orientation of the TEM lamella, we could observe that, in our case, some regions of this 2D layer are crystalline (Fig. 7 (c)). These regions might correspond to the formation of new NWs. It is also possible that the 2D layer is fully polycrystalline, although this could not be checked in this crosssection analysis. activated by a morphological defect on the silica surface. The NW growth is quasi perpendicular to the silica surface. Small deviation from the verticality, as observed in fig. 7 (a), might be due to local roughness of the silica surface, at the nanometer scale. Stoica et al. 6 have already mentioned the roughness of the silica surface as a possible cause of deviation from verticality. We investigate this hypothesis by comparing the roughness of fused and thermal silica, for which the widths of the angular distributions of the NWs differ by more than 3° (Fig. 5). KOH etching on these NWs (Fig. 9b) results in a flat end at the base and a sharp tip at the top, which confirms that the Ga polar facet is indeed at the base of the GaN NW. This strongly suggests that our NWs grow on silica with a N-polar top facet. The possible presence of inversion domains, which cannot be totally ruled out at this stage, will be discussed later. Recall that such N polarity is frequently observed in NWs grown by PAMBE on Si (111), 6H-SiC or AlN buffer layers 15,16,17,18 .

Bushes

For growth of GaN NWs on thermal silica sample, most of the time there are 'bush' like NW structures seen among the high density vertical NWs. The density of these bushes are around 6 x 10 6 cm -2 . The NWs on these bushes are strongly tilted and usually taller and wider in diameter compared to the vertical NWs on silica sample, indicating that bush NWs are formed much earlier than the normal vertical NWs. This let us to suspect that there could be something on the surface of the thermal silica layer which induces the formation of these NWs quite early. So we elaborated a sample with a short growth duration (120 minutes) and performed energy dispersive x-ray analysis on normal NWs and on bushes. No particularly different peaks were detected for the bush structures. However, by SEM images, we observed spots of lighter contrast in some places and on EDS analysis, we found that these spots accumulate much more Ga compared to the surroundings and also the peak for Cu presence is observed. As we have not found similar features (bushes or contrasted spots) for NWs grown on fused silica sample, we suspect that Cu contamination might be from the oxidation ovens used to obtain thermal silica. Though the reason for appearance of these bush structures is still not clear.

Optical quality

The optical properties of NWs grown on fused silica and on Si (111) were compared by analysing photoluminescence (PL) spectra of single NWs. The NWs were detached from their original substrate, dispersed by sonication in ethanol and deposited with a low density on a new Si/SiO2 substrate. A total of 6 and 5 single objects were analysed for the fused silica and Si (111) samples, respectively. Micro-PL spectra were measured at 4 K using the HR460 spectrometer with a 1800 grooves.mm -1 grating and a CCD camera. The energy resolution of the setup during these experiments was ~0.3 meV. Fig. 11 shows typical low temperature micro-PL spectra of NWs grown on Si (111) and fused silica. Note that these spectra may originate from single NWs, NWs having coalesced during growth or NW bunches aggregated during the dispersion 19 . In the NW-on-silica spectrum, the same near-band-edge emissions dominate, with however the FXA line stronger than the D 0 XA line. The FWHM of the D 0 XA and XA peaks are 4.27 meV and 12.7 meV for Si (111) and 17.6 meV and 10.4 meV for fused silica. The emission is broader for NWs grown on fused silica. The spectral components exemplified in fig. 11 were found for all investigated NWs, but the relative intensities of the bound and free exciton lines vary from spectrum to spectrum, possibly due to different local excitation conditions of the objects being measured (single NW, coalesced NW or bunch). Nevertheless, the free exciton line is always present. Its high intensity can be attributed to the relatively high excitation power used in our experiment 20 . In addition, other sub-bandgap contributions are observed in the PL spectra: a weak contribution near 3.42 eV is present for both samples and could be due to excitons bound to planar defects such as stacking faults 21 . A stronger contribution is detected at 3.45-3.46 eV for the NWs grown on Si (111). This contribution is much weaker, if not absent, for the NWs grown on fused silica and macro-PL experiments on as-grown samples confirm this tendency.

Based on the studies reported in chapter 2 The origin of IDBs is still a matter of debate 22 . Our upcoming reasoning about the lack of IDBs for NWs grown on silica is based on the two differences with respect to the growth on Si(111):

(
-Lack of a crystalline surface underneath: For GaN NWs grown on Si(111), AlN buffer is used. Thus the NWs are grown on a crystalline layer compared to the amorphous nature of the silica substrates. -2D layers between the NWs: Between NWs, GaN 2-D layers of Ga polarity is found commonly for growth on Si(111) with AlN pre-deposition. For silica samples, we observe comparatively a very thin (~5 nm), polycrystalline, continuous 2D layer is observed.

In GaN NW growth on Si(111) with AlN pre-deposition, it is known that Ga polar pyramids nucleate from the Al-polar regions. Due to energy constraints, this Ga polar pyramids do not grow in columnar form but rather in the form of a 2D layer. In the thesis of T. Auzelle 22 , it is proposed that this GaN polar pyramids when reaching the AlN regison suitable for N-polar GaN nucleation, gets surrounded by N-polar GaN crystallites and proceed to form NWs with IDBs.

We believe that the absence of IDBs in silica could be reasoned with this explanation. With the amorphous silica surface, we can safely say that there is no 'epitaxial driving force' to induce nuclei of different polarity. Stable nuclei of single polarity (N-polar) could be formed naturally due to the accumulation of significant volume of Ga and N adatoms. The point can also be supported by the fact that there is absence of a thick 2D layer between the NWs, which could confirm the absence of Ga polar crystallites in a non-epitaxial growth.

These points will be discussed for NW growth on graphene as well.

Conclusion and future scope

To conclude, we have shown that high quality GaN NWs can be grown on thermal or fused silica, using RF plasma molecular beam epitaxy. The NWs are N-polar, they have a random in-plane orientation but their axial orientation is remarkably aligned with the substrate normal. The average deviation from the normal orientation is as low as 2° for the fused silica substrate. This alignment could be related to the smoothness of the silica surface on which nucleation takes place. No extended defects are present at the GaN-silica interface.

Moreover, PL and TEM studies indicate that IDBs are less likely to form in NWs grown on fused silica than on Si (111). The incubation time which precedes NW formation is much longer on amorphous silica than on crystalline Si. This probably reflects the weaker bonding energy of the GaN NWs on the amorphous silica. This weaker bonding should be very well suited to an eventual transfer process: peeling off the NWs from the silica must be easier than from any crystalline substrate. The growth of near vertical GaN NWs on silica without degradation of their structural and optical properties opens new routes for the growth of GaN NWs on foreign substrates, provided the material can resist the GaN NW growth temperature. Therefore, growth on fused silica substrates could be used without restriction for the fabrication of NW-based optoelectronic devices.

A brief introduction to graphene

Graphene is an attractive candidate for a substrate material due to many of the following reasons: It is the thinnest crystalline layer possible, transparent for most of the wavelengths, has high mechanical flexibility, high thermal conductance (5.3 x 10 8 W/mK) 1 and has a high electron mobility (>40k cm 2 /V.s for monolayer graphene on silica) 2 . Additionally, it can be synthesized at low cost on large areas (as large as hundreds of metres 3 ). Furthermore, graphene films are transferable on almost any carrier substrate, including amorphous and/or flexible materials. This transfer can be performed before or after the deposition of an epitaxial material on the graphene surface. In fact, graphene is being used as a medium for easy peel off of epitaxial material from the support 4 . Even in cases where devices are fabricated on the as-grown bulk crystalline wafers, graphene is used a contact layer 5 . Therefore, a direct growth of semiconductor materials on graphene could open exciting perspectives for the fabrication of semiconductor devices without the need for a bulk crystalline substrate, where graphene will act as a transparent contact, epitaxial seed layer and a flexible carrier substrate.

Van der Waals epitaxy

In conventional growth of materials by MBE on bulk crystalline substrates like Si, the epitaxial material forms covalent bonds with the substrate. Due to the covalent bond formation, the lattice misfit produces crystal defects at the hetero-interface which reduces the performance of the devices fabricated from the epitaxial layers. In 1984, Koma et al. first demonstrated the growth of materials using non-covalent hetero-epitaxy and termed it as van der Waals (vdW) epitaxy 7 . On a surface without the presence of dangling bonds, the van der Waals forces between this passive surface and the precursors, drive the formation of epitaxial material. Though the exact nature of bonding at the interface between a covalent crystal and a 2D material such as graphene is a matter of debate, it is not a usual covalent bonding. For this reason, it is often referred as vdW epitaxy 6 . VdW epitaxy can produce epitaxial layers with clean interfaces 8 and well-defined epitaxial relationship 9 . However, as the lattice mismatch with the substrate can be accommodated through the non-covalent interface, vdW epitaxy leads to the formation of single-crystals with defect-free interfaces 8,9 . Thus, if the principle of vdW epitaxy applies to the growth of GaN on a 2D material like graphene, it is of great interest to obtain high quality GaN material.

In general, vdW epitaxy could denote on obtaining three kind of epitaxial material stacking order:

-2D material growth on 2D substrates -2D material growth on 3D substrates -3D material growth on 2D substrates

The focus of our work will be on obtaining GaN nanowires (3D material) on graphene (2D substrate).

State of the art: Graphitic substrates

Attempts to implement VdW epitaxy have been reported in detail for the cases of ZnO 12 , GaAs 13 and InAs 14 NWs. Regarding GaN, there have been attempts to grow epilayers on graphite 15 or graphene 16 , without or with intermediate material to

promote GaN nucleation, such as ZnO nanowalls 17,4 or

AlN nucleation layer 18 . The intermediate material is required as the GaN nucleation rate on pristine and defect-free graphene is reported to be low 12,16,17,18 . 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 17 . More specifically, growth of GaN NWs on graphitic surfaces was demonstrated by several groups. However the NWs grow along several directions 19,20 , unless a buffer layer is used 21 or unless graphene lies on a crystalline substrate 11 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. Recently the same group has used graphene as a buffer layer for growth of vertical GaN NWs on crystalline substrates such as Si(100 and Si(111) 10 . For comparison, this group transferred graphene on SiO2 and obtained in that case irregular structures with rather random growth directions 11 . These results reveal an important impact of the underlying substrate on the morphology and orientation of GaN crystallites. This suggests that, in this study, graphene could not be directly involved in the epitaxial alignment of the GaN structures; the epitaxial 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.

Objective

In the present work, we want to study in depth the growth of GaN NWs by PAMBE on graphene substrate placed on a non-crystalline support (SiO2). With such a stacking, we can exclude a possible in-plane alignment of the GaN structures induced by the carrier substrate. We used mono-layered and few-layered graphene obtained by chemical vapour deposition process. GaN NW nucleation dependence on the different number of monolayers is studied. Finally, we would aim to obtain selective area growth of NWs on graphene patches obtained by e-beam patterning.

GaN nanowires on graphene

Different type of graphene layers

Graphene layers were grown by chemical vapour deposition (CVD) on copper substrates or nickel layers.

By controlling the growth conditions, different types of graphene layers can be obtained. In this thesis, we used four different types of graphene layers:

-1 cm2 patch MLG (monolayer graphene): 1 cm by 1 cm patch of polycrystalline graphene layer with domain sizes in the range of µm 2 . This type of graphene comprises only one monolayer of graphene.

-1 cm2 patch FLG (few-layered graphene): 1 cm by 1 cm patch of polycrystalline graphene layer with domain sizes in the range of sq.µm. This type of graphene comprises typically 4-5 monolayers of graphene.

-Small graphene flakes: These are monocrystalline, isolated graphene domains. These flakes are around 600 nm in width and consists of only one monolayer of graphene. -Large graphene flakes: Typical dimensions of these kind of flakes are in the order of tens of µm. Each flake is an isolated monocrystalline domain having one monolayer at the periphery and additional monolayers (up to 5-6 monolayers) are stacked near the centre of the flake. This stacking can be either in the form of a 'wedding-cake' structure or 'inverted wedding cake' structure 24 .

The morphology of such multi-layered graphene flakes is often reported and has been studied in detail for similar growth conditions 25 . It is important to note that the edges of the graphene flakes grown on copper foils is well established to consist of carbon zigzag termination 26,27 .

Preparation of graphene patch samples

Mono-layered and few-layered graphene were procured from Graphene Supermarket. These graphene patches were grown by chemical vapour deposition (CVD) on copper foils and transferred on a silicon substrate with an oxide layer of 500 nm and a nickel layer of 100 nm, respectively for shipping. Details about the CVD growth of graphene can be found in literature 28,29 . The number of layers in few-layered graphene has been characterized by separate Raman and atomic force measurements. The average thickness is found to be equal to 4±1 monolayers. 1 cm 2 pieces of monolayer and few-layered graphene were transferred by a wet transfer method to a silicon (001) substrate with a nominal 25 nm SiO2 layer formed by dry oxidation. To transfer monolayer graphene, the graphene on copper foil was first covered with 50 nm-thick polymethyl methacrylate (PMMA) protective layer, and the copper foil was etched using (NH4)2S2O8 solution. The floating graphene covered by PMMA was transferred to deionized water, then deposited onto the SiO2/Si template, and the samples were naturally dried. The PMMA protective layer was removed by cleaning with solvent (acetone, isopropyl alcohol) and deionized water. To transfer few-layered graphene, the same PMMA protecting layer was used. The Ni layer was wet-etched by FeCl3-saturated aqueous solution, and the floating graphene sheet was transferred to clean deionized water and then deposited on the SiO2/Si template and naturally dried. It should be noted that contrary to what is seen in literature 30,31 buffered oxide etch (BOE) of SiO2 was not used to avoid modifying the graphene properties. The Ni layer was etched from the lateral surfaces for approximately 12 h.

Before the MBE growth, to minimize the water residues at the graphene/substrate interface for both few- layered and monolayer graphene, the samples were annealed at 130 °C using a slow temperature ramp on a hot plate in ambient conditions during 1 h.

Preparation of graphene flake samples

Monocrystalline graphene domains were home-grown at our laboratory by Dr. Ali Madouri. Graphene flakes were synthesized on copper foils by pulsed-chemical vapour deposition (CVD) technique. To obtain hexagonal graphene domains, the experimental conditions are a gas pressure of 25 mbar, a growth temperature of 1000 °C and the use of ten cycles of CH4 (20 sccm, 5 s) and H2 (960 sccm, 55 s). These conditions refer to the work of V. Bouchiat et al. 32 and were applied to our commercial Aixtron BlackMagic system. We can obtain isolated graphene domains which are mostly hexagonal, have one monolayer height and approximately 40 μm width. For larger domains, new graphene layers start to form at the center of the domains. The final morphology of a flake consists of several terraces forming a pyramidal (or wedding cake) structure. 1 cm 2 patch of isolated graphene domains was transferred by a wet transfer method to a Si (100) substrate with a nominally 300 nm SiO2 layer formed by dry oxidation. To perform this transfer, the graphene domains on copper foil were first spin-coated with a 350 nm-thick polymethyl methacrylate (PMMA) protective layer, and then dipped into (NH4)2S2O8 solution in order to etch the copper foil. The floating graphene covered by PMMA was transferred to deionized water and profusely rinsed, then deposited onto the SiO2/Si template, and the samples were naturally dried. The PMMA protective layer was removed by cleaning with solvents (acetone, isopropyl alcohol) and with deionized water.

Growth of GaN NWs

After the transfer of graphene on SiO2/Si(100) templates, the samples were loaded in the MBE chamber to grow GaN NWs. Neither buffer layer nor catalyst particles of any type were deposited before the 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 radiofrequency power of 350 W. During growth, the background pressure was 8.3x10 -6 Torr and the substrate temperature was maintained at 800 °C. 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 and motivates to try the growth on monocrystalline graphene flakes.

On graphene monolayer patch

At the edges of the graphene monolayer patch, (seen in Fig. 6) it can be seen that there is no more NW on the SiO2 layer surrounding it. This clearly shows that the NWs grew selectively on graphene and that their orientation is imposed by graphene and not by the support underneath.

On graphene flakes

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 mono-crystalline graphene flakes. After transferring the isolated graphene flakes onto a SiO2/Si template, GaN NW growth was performed under the same conditions as that of the previous sample. The image in fig. 7 (a) shows the top view SEM image of the NWs grown on a small graphene flake. In consistent with the previous observation, we found that the growth is again highly selective i.e., NWs are present only on the graphene flake and not on the SiO2 templates.

All the NWs are again vertical and their density is uniform across the flake. In the magnified SEM top-view (Fig. for the whole graphene flake since it has a monocrystalline structure.

The next growth experiment was performed under similar conditions but on a larger graphene flake having the wedding cake morphology as described before. Referring to detailed studies of similar few-layered graphene domains,<sup>25</sup> we can reasonably assume that the flake is pyramidal with monolayer steps, i.e. the number of graphene monolayers decreases by one unit every step from center to periphery. The NWs and will be discussed later in section 5.4.3.

Silica intermediate layer

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. In order to study the influence of thickness of SiO2 carrier layer on NW verticality, 1 ML graphene was transferred on two different samples with SiO2 thickness of 50 nm (sample i) and 300 nm (sample ii). Fig. 9 shows 45° tilted view and cross-sectional view of these two samples. Both the samples are grown at 800°C growth temperature, with a V/III ratio of 1.1 and growth time of 2 hours.

As can be seen from the SEM images, NWs have identical morphology, with good verticality, on both the samples. The morphology is also the same as in growth experiments on a 25 nm thick SiO2 layer described in the previous section. The slight differences in NW heights and densities are attributed to small variations in growth temperatures and molecular fluxes between two growth experiments. By top-view SEM observations, we verified that the NWs possess same in-plane alignment characteristics as seen before. Therefore, we conclude that the thickness of SiO2 does not have an influence on the alignment of GaN NWs.

Characterisation

XRD and Pole-figure results

X-ray diffraction (XRD) was used to determine the NW growth direction on the graphene monolayer patch sample (mentioned in section 5.2.2.1). A 2 theta/omega analysis (Fig. 10 (a)) 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. As usually observed, the GaN NWs grow along their c-axis. We quantified the verticality of the NW ensemble by recording an XRD rocking

TEM results

We examined the crystalline structure at the NW base using transmission electron microscope analysis (TEM). The TEM observations as seen in Fig. 11 show a perfect wurtzite structure, free of defects and start vertical right from the base. Our TEM analysis could not reveal the atomic arrangement of the graphene layers below the NWs.

There are many reasons such as:

-Damage of the graphene layer by the 200 keV e-beam in the TEM; This is unlikely as graphene layers up to 1 ML have been spotted in our laboratory using the same setup and parameters.

-Long exposure to the Nitrogen plasma during growth; This is again unlikely as at least the graphene under NWs should be protected as we have the in-plane orientation of the NWs.

-After due consideration, we tentatively associated the lack of observation of graphene to the surface roughness of the SiO2 layer underneath.

These reasoning could be verified by further experiments like:

-Observation of graphene by Raman spectroscopy before and after exposure of either e-beam or long exposure of nitrogen plasma.

-TEM observation of GaN NWs grown on graphene placed on atomically flat SiC wafer. These experiments are planned to continued further in our laboratory.

Photoluminescence results

The optical properties of the NW ensemble obtained on the monolayer graphene sample were studied by photoluminescence (PL) spectroscopy. Fig. 12 shows the spectrum collected from a NW array grown on a 1 cm 2 patch of 1 ML thick graphene (described in 5.2.2.1). PL for GaN NWs on graphene domains was also measured and the spectral response was the same. 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 eV, 3.45 eV and 3.42 eV with FWHMs of 14 meV, 24 meV and 31 meV respectively. The 3.472 eV peak (D 0 XA) is relatively wide, but comparable to that we obtained for GaN NWs grown on conventional Si (111) substrates 33 . The 3.45 eV peak (attributed to IDBs) 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) 33 , a possible sign that NWs grown on graphene contain less inversion domain boundaries than those grown on Si (111). In continuation with the discussion mentioned in chapter 4, we can denote again that the reduced intensity of 3.45 eV emission, could be related to the lack of strict epitaxial driving force to form a Ga-polar crystal and hence the IDBs. This first thought about IDB formation will be examined with more experiment in the future. 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 34,35 . Further investigations will be carried out in our laboratory to verify this attribution. A similar peak is also observed in our GaN NWs grown on Si (111) substrates 33 . Therefore, the optical quality of GaN NWs grown on graphene-on-amorphous templates is not degraded with respect to the GaN NWs on a crystalline Si substrate.

Epitaxial growth of GaN on graphene

Lattice relationship with graphene

As we have shown that GaN NWs seem to have a defined relationship with graphene, we wanted to further investigate this topic. First, we analysed the relative in-plane orientation of the two crystal lattices by simple geometrical observations by SEM. We check this relative orientation by diffraction analysis. Finally, we propose possible atomic arrangements which satisfy the relative in-plane orientation and which seem plausible with respect to a strain sharing model. We determined the orientation of the NW facets relative to the edge of the graphene flakes (Fig. 13).

As seen in the previous sections, all the NWs grown on a single flake have the same in-plane orientation, their lateral facets corresponding to the {101 ̅ 0} planes 36,37 . Schematics of the graphene flake of hexagonal shape is shown in the inset of Fig. 13a as a black hexagon. The detail of the honeycomb lattice of C atoms at one corner of the flake is shown in Fig. 13c. The edges of the hexagonal flake are of zigzag type, as known from the literature on such graphene flakes 26,27 . The yellow hexagon in the inset of the honeycomb carbon lattice. This in-plane relationship is found to be the same for the GaN NWs grown on the larger monocrystalline flake, with up to five staggered monolayers at the centre of the flake. Another NW sample grown on a wedding cake graphene flake is further investigated by selective area electron diffraction in a transmission electron microscope. We want to check whether the NW in-plane orientation depends on the number of graphene MLs underneath.

SAED patterns

To unambiguously determine the 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 6 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 (Fig. 14 ensembles from each zone (Fig. 14 (c)). 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 (Fig. 4 (a)). The in-plane orientation of the first graphene layer is thus preserved when the pyramidal stack forms. Though the SAED patterns does not cover the 1 ML step of the flake, by top-view SEM observation, we have also checked that that the NWs have same inplane orientation as that of the NWs on the other steps. Some SAED patterns present curved spots indicative of some deviation in the vertical alignment of the NWs. This might result from the non-ideal 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 [11 ̅ 00] direction is parallel to the long axis of the lamella, i.e. 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 (Fig. 13). 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. 13 . According to their prediction, the m-plane facets of NWs would be aligned with a rotation of +/-19.1 o to the carbon zig-zag edges based on misfit strain a. . This second possibility is less likely as in a unit cell, lesser number of atoms sit on the energetically favourable site compared to the number of atoms that do not.

Epitaxial growth configuration

Difference in NW nucleation with respect to different number of MLs of graphene*

* I want to acknowledge Frank Glas' contribution to this part. In particular, he developed the model based on elasticity which provides a plausible explanation to the experimental observation.

I will discuss now about 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 39,40 which show the formation of Au nanostructures with a density decreasing with increasing number of graphene layers 39,40 . 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 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. Referring to fig. 8, it is clear that the volume of each NW and the NW density decreases simultaneously with respect to the number of layers. The total quantity of GaN material deposited per unit area is estimated to vary from 2e 8 nm 3 .µm -2 on 1 ML of graphene to 6e 6 nm 3 .µm -2 on 5 ML of graphene.

Another explanation could result from the Ehrlich-Schwoebel barrier associated to 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 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 was only to assess the magnitude of this effect and its possible influence on nucleation and hence on NW density. To this end, we adopted a very simple model. We assumed flat 2D nuclei and a compliant graphene flake (no strain is induced by the carrier substrate) and neglect strain relaxation at the nucleus edges. As regards 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:

𝑤 𝑒 = 𝑛 𝑛 + 1 𝐸 1 -𝜈 𝜖 0 2 ( 1
)
with 𝜖 0 the in-plane misfit between graphene and GaN and n the number of graphene layers. Using Equation is the same on all terraces. Figure 18 shows that Equation (2) fits the measured n-dependence of the NW density very well. From this fit, we extract the 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.

Figure 18: 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 8b) using eq 2. A continuous curve is given as a guide, although the density is only defined for integer values of n.

Growth on micro-and nano-patches of graphene ML

NWs obtained on the micro-sized graphene patterns are highly vertical and are highly dense as in the case of self-aligned NWs on graphene monolayers. When the distance between the edges of two patterns is less than or equal to 1 µm (pitch of 5.4 µm and pattern width of 4.98 µm), the selectivity is quite high (97.42%) (as seen in Fig. 20 (a)) with almost no NWs on the silica layer. Selectivity is measured as ratio of density of NWs obtained on graphene to the sum of densities of NWs obtained on graphene and silica. No bushes were found on silica for this pitch. For larger inter-pattern distances, we observed bush like NWs on the silica layer. We have seen such kind of bushes for NWs grown directly on silica and we associated this to some surface contamination present in the thermal silica or introduced during the graphene transfer process. A particular feature of these bushes is that they are never present closer than 1µm to the graphene patterns. This shows that contamination is not the only reason for having these bushes.

The unwanted NWs related to surface contamination are more numerous on the nano-patterned samples, possibly due to the larger distance between the patterns and their reduced size. Apart from the NWs due to the surface contamination, we have achieved 100% selectivity for growth on graphene patterns. This is achieved by using high growth temperatures (~815 o C), which inhibits the nucleation of GaN NWs on silica.

Though we are still far from having a perfect selective growth due to the surface contamination, with these results we have provided a proof-of-concept that it is possible to use patterned graphene layers to obtain SAG of GaN NWs on silica supports by PAMBE. The optimization of SAG is continued in the laboratory. Our end goal will be to obtain 1 NW per graphene pattern.

Conclusion and future scope

In conclusion, we have identified that CVD-grown graphene is an excellent substrate for GaN NW growth by PA-MBE. Epitaxy is demonstrated on large and small graphene flakes transferred onto amorphous SiOx carrier layers. In our growth conditions, NW growth is highly selective; NWs nucleate only on graphene layers and not on the surrounding silica. The NWs grow vertically along their c-axis and we observed a unique epitaxial relationship with the 〈211 ̅̅̅̅ 0〉 directions of the wurtzite GaN lattice parallel to the directions of the carbon zigzag chains. We proposed that the two materials are strained to accommodate a -3.1% in-plane misfit insuring a certain lattice coincidence. Strikingly, the number of graphene layers had 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 which, at NW nucleation, partitions between graphene and GaN nucleus. The higher number of graphene layers, the larger is the elastic energy stored in the system, which elevates the nucleation barrier. 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 reference 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 and for achieving NW organized arrays with a high control of the wire-to-wire homogeneity.

Conclusion

To conclude we studied three different approaches to grow GaN NWs on amorphous supports. In a first approach, we started from the amorphous substrate (SiOx) and formed thin films (6 nm) of polycrystalline Si by aluminium-induced crystallization. The conditions of this process were optimized in order to get a strong [111] fiber-texture of the Si film which enabled us to grow vertically oriented GaN NWs. Combined with lithography, we used this process to produce organized arrays of mono-crystalline Si nano-platelets, each of them representing a nano-substrate. On these nano-substrates, we achieved selective area growth of NWs. It shows that a precise control of the NW nucleation sites is possible. However, the high growth temperature (> 810 o C) of GaN was found to alter the morphology and the crystalline characteristics of the Si platelets. This motivated us to explore a second approach.

In the second approach, the possibility of growing NWs directly on amorphous substrates (thermal and fused silica) was also explored. Self-induced GaN NWs were formed with a good verticality on both substrates, i.e. without the need of crystalline order to direct the growth. We associated this high verticality to the smooth surface of the silica layers as indicated in literature. TEM analysis show that NWs does not have any defects at the base and start vertical right from the substrate. In terms of optical quality, the NWs have almost no inversion domain boundaries in them.

As a final approach, graphene was used as a substrate on top of an amorphous silica support. Three different types of graphene layers were used: cm 2 patches of mono-layered and multi-layered graphene, smaller (nanometre range) and larger (micrometre range) monocrystalline domains of pyramidal shape and finally patterned domains. We obtained dense arrays of vertical NWs on the graphene areas with, again, a very good selectivity (no growth on silica surface). The NWs and the graphene layer have a single relative in-plane orientation which we determined for the first time in literature. We proposed a plausible epitaxial relationship of the two crystalline lattices implying a -3.1% misfit. We demonstrated that the number of graphene layers has a strong impact on GaN nucleation: the NW density and height decrease rapidly when the number of graphene layer increases from 1 to 5. We proposed a model that explains quantitatively this striking phenomenon.

Patterned graphene nano-substrates are also fabricated by lithography and successfully used for localized NW growth. PL measurements show that there is reduced signal for IDBs as compared to the NWs on Si(111).

As a general conclusion of these investigations, the epitaxial growth of GaN NWs on graphene offers a wide choice of possible underlying supports. With a growth temperature higher than 810 o C it is possible to obtain selective area growth which is an important aspect in case of device fabrication. As of now this is not feasible with AIC-Si (due to deformation) and silica substrate (lack of a suitable mask material). With graphene substrate, the proposed strain sharing mechanism could also help in obtaining NWs with reduced defects even at their base. Additionally, this approach looks particularly promising for the development of flexible devices where graphene could also act as a transparent electrical contact material for GaN NW devices.

INTRODUCTION

In this section a brief introduction is provided for different topics which could help the readers understand better the manuscript. As this thesis deals mainly with growth of gallium nitride, a brief introduction to the material properties and the defects is provided in section A.1. Section A.2 deals with the experimental setup of our molecular beam epitaxy used to grow GaN NWs and section A. Inversion domains are basically parts of the crystal where the polarity is different compared to the rest of the crystal. As the result the boundary between these two domains does not follow the usual bonding (alternating anion-cation bonding) and it is called the inversion-domain-boundary. For applications based on using the polarization effect of GaN, IDBs are detrimental as the inversion of polarity could mean in the inversion of voltage leading to the reduced output of usable power from the material.

IDBs usually arise from the substrate-film interface, however the exact reason for their occurrence is still under a matter of debate. By various reports, it has been attributed to the inhomogenous nitridation of the substrate 3,4 polarity of the buffer layers 5 , imperfect substrate preparation 6 , steps in the substrate 7 and buffer layer thickness 8 .

A.3 Plasma assisted molecular beam epitaxy (PAMBE)

The term "epitaxy" comes from the Greek terms "epi" and "taxis" meaning "above" and "orderly" respectively. The term "molecular beam" refers to the fact that the evaporated species contributing to the growth are characterized by a long mean free path and thus do not react with each other or with the gas in the chamber until they reach the substrate. Today MBE is a reference growth technique for the synthesis of monocrystalline and nanomaterials for high quality and highly controlled dimensions using the "bottom-up" approach. This technique is performed under high vacuum conditions (10 -6 to 10 -12 Torr). With the absence of carrier gas and the availability of high vacuum, the MBE achieves a material purity level unattainable by other epitaxial techniques. The relatively low growth temperatures can limit inter-diffusion phenomena at interfaces and cells containing individual species with shutters (opening and closing) makes it easy to change materials and hence the diffracted beam is oriented in all the directions. Thus the diffraction pattern is in the form of concentric circles. And finally if the surface is covered by monocrystalline islands, the diffraction of the electron beam on these structures results in the appearance of additional lines in the RHEED pattern. In particular, the blocks with a high aspect ratio, such as nanowires, induce normal periodicity on the surface. This out of plane

-

  Thin layer of polycrystalline silicon fabricated by Aluminium induced crystallised Silicon (AIC-Si) -Thermal and fused silica substrates -Graphene layers fabricated by chemical vapour deposition This will not only reduce the use of raw materials but also open a perspective of adding new optical, electrical or thermal functions to the support.

Figure 1 :

 1 Figure 1: A brief history of Gallium nitride crystal growth, between 1930 and 2000.

Figure 3 :

 3 Figure 3: Percentage lattice mismatch with GaN vs cost of a 2-inch wafer (in euros) for different substrates available in market. Lines connecting the points are just a guide to the eye. Triangle marked with a question mark could be the place of an ideal substate for GaN growth.

  locations for NW formation. Substrate patterning is also efficient to control the NW density. NW-based device fabrication is being studied in details by various groups, using both axial and radial heterostructures. From a geometrical point of view, radial heterostructures of high form factor provide larger active area for the same amount of substrate surface. Hence, for a given drive current in a NW-LED, the carrier density should be lower in this configuration and thus less sensitive to Auger recombination. Hu et al. have shown that NWs have better absorption at high frequency regime and by adjusting

Figure 4 :

 4 Figure 4: Critical dimensions (Height and radius) for NWs with variation of lattice misfit in percentage. Taken from Glas et al.27

Figure 5 :

 5 Figure 5: Illustration of a NW device with (a) axial junction (b) radial junction

4 .

 4 UHV chamber offers the advantage of using in-situ observation techniques such as reflection high-energy electron diffraction (RHEED) (explained in detail in Annex A.4) during growth. In plasma-assisted molecular beam epitaxy, Ga is provided by a Knudsen cell and N active species are supplied by a plasma cell. Beams of these atoms are focussed onto the hot substrate surface. The atoms supplied to the substrate surface undergo one or many of the following processes during the growth by MBE (as shown in the following figure): Incorporation to the crystal lattice The impinging atoms are adsorbed on the surface and can desorb or diffuse around and reach favourable sites of incorporation (existing surface steps or kinks) or form a cluster of several atoms on a flat terrace. These clusters are generally metastable nuclei (they tend to decompose) until a critical radius is reached at which, the nucleus becomes stable, i.e. its expansion is energetically favourable and a monolayer grows epitaxially with the substrate surface. Initiation of GaN NW growth on a foreign substrate also requires a nucleation stage which can last minutes or tens of minutes depending on the growth conditions and the nature of the substrate surface. The time taken since the supply of the fluxes to the formation of stable nuclei is called 'incubation time'.

Figure 6 :

 6 Figure 6: Illustration showing the mechanisms that can occur on the substrate: Adsorption, desorption, diffusion and incorporation on the crystal lattice

Figure 7 :Figure 8 :

 78 Figure 7: Three different growth modes of material growth on an epitaxial substrate

1 . 3 . 2 Figure 9 :

 1329 Figure 9: Illustration showing the concept of using thin polycrystalline films formed on amorphous substrates for vertical NW growth. Inspired from Cohin et al.32

Figure 10 :

 10 Figure 10: (a) 45° view SEM image showing vertical NWs grown on TiN film and (b) a close-view of the same showing that NW have a good morphology. (c) Photoluminescence measurements for GaN NWs grown on TiN show identical results as that grown on Si(111) substrate. Taken from Wolz et al.35

Figure 11 :

 11 Figure 11: (a) illustration showing the mask layer deposited on polycrystalline Si layer. SEM images showing (b) the overall results for growth on polycrystalline layer with formation of either (c) NWs (d) or hillocks (e) or no growth.Taken from Ikejiri et al.36

Figure 12 :

 12 Figure 12: Illustration showing (a) Ti pre-orienting layer on glass followed by a fiber textured GaN nucleation layer and (b) the overall scheme of SAG of single crystalline GaN pyramids grown on the GaN nucleation layer. Taken from Choi et al.37

Figure 1 :

 1 Figure 1: Variation of GaN NW growth with respect to growth temperature using (a) using Al first AlN predeposition and (b) N-first AlN pre-deposition. Taken from Auzelle et al.4

4

 4 Figure 1: Variation of GaN NW growth with respect to growth temperature using (a) using Al first AlN predeposition and (b) N-first AlN pre-deposition. Taken from Auzelle et al.4

Figure 2 :

 2 Figure 2: (a) Top-view SEM images of GaN NWs grown on different thicknesses of AlN (b) Electroluminescence of GaN NWs grown on AlN layer and on bare Si layer. Taken from Musolino et al. 5

Figure 3 : 1 ) 6 ) 5% HF for 1 minute 7 )

 3167 Figure 3: HAADF STEM image of a IDB-free (#NW1) and IDB-present (#NW2). The µ-PL for these NWs showing 3.45 eV emission only in #NW2. Taken from Auzelle et al. 6

7 x 7

 7 fig.4.

Figure 4 :

 4 Figure 4: Illustration showing the Si (111) surface reconstruction

Figure 5 :

 5 Figure 5: Illustration showing the RHEED intensity variation with respect to the stage of a monolayer growth

  AlN pre-deposition, starting with Al alone, before the growth of NWs to prevent the formation of SiNx. This is our standard procedure.To deposit Al, after the flux calibration, Si (111) substrate temperature is stabilized at 620 o C and the surface is exposed to an Al flux of the order of 8.10 -8 Torr for 40 seconds. Normally for this flux and duration, we deposit 6 MLs of Al. With this thickness of Al, Si (111) surface is partly covered with Al (so called the γ-phase) and the remaining excess Al is accumulated in the form of droplets12 . The aluminium droplets form an Al-Si eutectic composition with the substrate. This pre-deposition of Al is followed by the ignition of the nitrogen cell and exposure of the surface to N species. For nitridation process, the substrate is maintained at 620 o C and is exposed for 60 seconds to active N flux equivalent to 0.7 ML.s -1 . Between the Al deposition and the N exposure, the distribution of aluminium droplets evolves by Ostwald ripening, which affects the final structure of the layer (and consequently the NW nucleation). Hence for reproducible results, this switching duration must be maintained the same for all the samples and it is typically between 90-120 seconds for our samples. Upon Al exposure to nitrogen flux, several phenomena can occur simultaneously: -Formation of incomplete ML of AlN: The γphase corresponds to a reconstruction of the Si dangling bonds of the surface involving the Al atoms. This surface is converted into an incomplete monolayer of AlN and is of Al polarity. -Formation of AlN nanocrystals: Al-Si eutectic droplets react with the N radicals and result in growth of Si doped nanocrystals of AlN. It has been shown by Largeau et al. that each

Figure 6 :Figure 7 :

 67 Figure 6: Illustration showing the procedure to grow GaN NWs on Si (111) without and with AlN pre-deposition

Figure 8 :

 8 Figure 8: RHEED pattern corresponding to the (a) AlN deposited (before NW nucleation) and (b) after NW nucleation.

Figure 9 :

 9 Figure 9: Growth diagram for our MBE chamber. Red dots represent each sample grown at that particular temperature and V/III flux ratio.

Figure 10 :

 10 Figure 10: SEM images taken in (a) Top view showing high density GaN NWs, (b) Cross-view to showing that the NWs are highly vertical and (c) high magnification top-view showing hexagonal facets of NWs.

Figure 11 :

 11 Figure 11: (a) pole-figure taken for GaN NWs grown on Si (111) with AlN pre-deposition indicating that the NWs are in-plane oriented (b) NWs have high-verticality (FWHM=0.5°)

Figure 12 :

 12 Figure 12: Graphic representation using the ball-and-stick model for Ga-polar and N-polar termination in GaN.

Figure 13 :

 13 Figure 13: Cross-view SEM images of GaN NWs as grown and after KOH etching

Figure 14 :

 14 Figure 14: Low temperature (4K) µ-PL spectrum taken for GaN NWs grown on Si (111) with AlN pre-deposition

Figure 15 :

 15 Figure 15: (a) Dark field and (b) bright field STEM images of GaN NWs grown on Si (111) showing the presence of IDBs in them. NWs on (a) are observed from {112 ̅ 0} zone axis

  in producing a module. An overview of the process is illustrated in fig. 1. Typically for AIC-Si, a stack consisting of a thin layer of amorphous silicon (a-Si) and a thin layer of aluminium is deposited on an amorphous support. Upon annealing at a temperature less than the eutectic temperature, the layers interdiffuse and eventually exchange resulting in the formation of a polycrystalline Si (pc-Si) layer. By this method, Si films with thicknesses ranging from nanometers to micrometers could be produced. The process fits well within the framework of developing a low-cost substrate for GaN NWs considering the following points: (i) Low-temperature process (ii) which means a cost-efficient support (glass) (iii) could use deposition techniques already available in industry (sputtering and annealing) and (iv) relatively short crystallisation time (~2 hours) thereby providing higher throughput.3.1.3 Potential substrate for NW growth 3.1.3.1 Fiber texture The resulting Si thin film is polycrystalline comprising domains oriented in random directions. However, in most cases one preferential direction of the Si domains is aligned perpendicularly to the surface of the support (will be seen in detail in section 3.3.1). This feature of having the common orientation of a single family of planes though the other planes are randomly oriented with respect

Figure 1 :

 1 Figure 1: An illustration of a typical AIC-Si process. On annealing, a stack of Al/a-Si layers, polycrystalline Si (pc-Si) is formed with the interchange of the layers.

Figure 2 :

 2 Figure 2: Illustration of a fiber-textured layer. Random inplane orientation with a common vertical plane

Figure 3 :

 3 Figure 3: First results for growth of NWs on AIC templates (a) 45 o view of GaAs NWs on an in-continuous AIC-Si layer (b) 45 o view of a perfectly vertical GaAs NW showing proof for possiblity of vertical growth (c) cross-view GaAs NWs on AIC-Si layer showing high number of vertical NWs.Image is taken from Cohin et al.24

Figure 4 :

 4 Figure 4: A method to grow Ge NWs on flexible plastic support by use of AIC-Ge layer (a) Schematics of the method to fabricate NWs on flexible supports without a transfer (b) A picture showing the flexibility of the substrate (c) 45 o view SEM images of AIC-Ge layer before NW growth and vertically grown Ge NWs on them. Image is taken from Toko et al.25

Figure 5 :

 5 Figure 5: (a) 45 o view SEM image of vertical Si NWs growing on patches of on AIC-Si template (b) ESBD pattern for AIC-Si template before growth showing only patches of AIC-Si layer have the 111 fiber-texture. Image is taken from Kendrick et al.28

Figure 6 :

 6 Figure 6: Cross-sectional SEM view showing the contact angle formed by the droplets when (a) there is no surface oxide (b) thicker surface oxide (c) and optimal surface oxide. (d),(e) and (f) shows the illustration of the end result based on the contact angle formed by the droplets. Image is taken from Ren et al.29

Figure 7 :

 7 Figure 7: Illustration showing the mechanism involved in AIC-Si process (a) Showing the interdiffusion of atoms into each other layer (b) Initiation of crystalline Si (pc-Si) nuclei in Al layer (c) Expansion of pc-Si nuclei (d) Exchange of layers and formation of complete pc-Si layer

  into the Al layer and formation of c-Si, creates a stress gradient with Al layer having compressive stress and the a-Si layer undergoing a tensile stress. As a result of this, the Al atoms from the Al layer starts to move towards the volume-free places of the a-Si layer. This results in complete replacement of a-Si layer by the Al layer and the formation of pc-Si at the place where Al layer was previously present, resulting in the AIC stack exchange. 3.2.1.1 Al/Si ratio As the a-Si layer wets the Al grains, c-Si nucleates between the grains and this crystalline phase grows. The layer thicknesses of Al and a-Si play an important role in the final pc-Si layer morphology. If the thickness of a-Si is less than the thickness of Al, there is insufficient material to form a continuous pc-Si layer. As a result, the final pc-Si is in the form of islands of crystal. If the thickness ratio is approximately equal to 1, the pc-Si forms a continuous layer by replacing the equal amount of Al. On the other hand, if the

Figure 8 :

 8 Figure 8: Illustration showing the three scenarios of AIC-Si formation w.r.t Al/Si ratio (a) Thickness of Si (hSi) less than the thickness of Al (hAl) leads to the formation of pc-Si island (b) Al/Si ratio equal to one gives an uniform layer of pc-Si (c) Al/Si ratio less than 1 gives a continous pc-Si layer and formation of pc-Si islands on top of the layer

3. 2 . 2 . 1 . 1

 2211 Deposition of the layers There are different means of depositing thin film of amorphous semiconductors. Chemical vapour deposition (CVD) techniques like low-pressure CVD (LPCVD), PECVD, MOCVD, APCVD can be used to prepare amorphous semiconductors at low temperatures. Other techniques include sputtering (DC sputtering, DC magnetron sputtering, RF sputtering), evaporation (thermal evaporation, ebeam evaporation) and reactive PVD. Though PECVD is the most common method in industry, in our experiments, we have used DC magnetron sputtering for the synthesis of thin layers of a-Si and Al, mainly for the relative low cost and high uniformity of the deposited layers. As seen before, a typical AIC-Si stack layer consists of layer of Al, thin oxide layer and a layer of a-Si. In most cases, the intermediate oxide layer is prepared by exposing the Al layer to ambient conditions overnight (few hours) to form native AlOx and then depositing of a-Si layer on top of it 34,35 . This is not only time consuming but also produces a completely uncontrolled interfacial layer of SiO2. In our case we used an in-situ oxide preparation method to produce a controlled layer of SiO2 intermediate layer. IN-SITU, CONTROLLED SiO2 LAYERS By standing at a constant power of 100 W, the potential of a Si target changes over a wide range when the rate of O2 in the gas mixture is modified. In the absence of O2, the target is metallic and acquires a potential of 474 V. When the rate of O2 increases (controlled by the SCCM flow), the potential of the target could increase up to 503 V, before rapidly decreasing to 319 V. The last plateau corresponds usually to the stoichiometric SiO2.

Figure 9 :

 9 Figure 9: Graph showing the fiber-texture orientations shown by different research groups as a function of IMO thickness (in terms of air exposure time) and thickness of Al (nm). It can be seen that absence of IMO layer leads to random orientation of crystal.

Figure 10 :. 4 . 3 . 2 . 2 . 1 . 2

 10432212 Figure 10: Voltage variation of the Si target with respect to the O2 flow in the chamber. It has been shown in the previous thesis of Y. Cohin that this variation is linked to the formation of surface oxide on the Si target

  Figure 13: SEM image and the corresponding illustration for the formation of pc-Si layer by varying the thickness of a-Si layer with a 10 nm Al layer. Pc-Si layer is obtained in the form of (a) islands for 5 nm a-Si (b) non-continuous layer for 10 nm a-Si and (c) islands + non-continuous layer for 15 nm a-Si.

Figure 14 :

 14 Figure 14: Top-view SEM images of pc-Si layer formed with a 15nm a-Si and 6nm-Al stack leading to a continuous layer formation

Figure 15 :

 15 Figure 15: Profile of the pc-Si layers formed by 6nm Al stack (6 nm a-Si/15nm Al) and standard stack (15 nm a-Si/15nm Al). 6nm stack has a smoother surface with the absence of holes to the second level of MIC.

Figure 16 :

 16 Figure 16: GIXRD 2theta-chi measurements taken for 6nm Al (6 nm a-Si/15nm Al) stack and standard stack (15 nm a-Si/15nm Al) showing a reduced <111> fiber-texture for 6nm Al stack.

Figure 17 :

 17 Figure 17: (a) Top-view SEM images for pc-Si layer formed by varying the composition of oxide layer by varying the SCCM flow (b) GIXRD 2theta-chi analysis showing an improved fiber-texture for thinner IMO for 6nm Al stack (15 nm a-Si/6nm Al)

Figure 18 :

 18 Figure 18: Illustration showing the steps involved in GaN NW growth on AIC-Si templates using AlN pre-deposition step.

Figure 19 :

 19 Figure 19: Top-view SEM images of GaN NWs grown on AIC-Si templates (a) with AlN pre-deposition and (b) without AlN pre-deposition.

Figure 20 :

 20 Figure 20: Pole-figure measurements taken for GaN NWs grown with (a) AlN pre-deposition showing broad deviation of NWs (b) without AlN pre-deposition showing high verticality. (c) Cut-line analysis provides a quantitative comparison of the deviation of NWs.

Figure 21 :

 21 Figure 21: (a) Pole-figure measurement for a pc-Si layer formed by AIC-Si template. The low brightness ring at 71° shows the presence of <111> fiber-texture in the pc-Si layer. (b) Cut-line analysis shows that the <111> fiber-textured grains are deviated by +/-of 6° to the normal of substrate.

Figure 22 :

 22 Figure 22: 45° SEM view showing a pc-Si layer (in the form of islands) (a) before growth of GaN NWs (b) after growth of NWs without AlN pre-deposition and (c) with AlN pre-deposition. (d) Top-view SEM image of a Al-Si pedestal possibly formed due to the melt-back etching of Si by Al while deposition of Al at 620 °C.

Figure 23 :

 23 Figure 23: Cross-View SEM images of GaN NWs (a) without and (b) with AlN pre-deposition after KOH etching showing pointed tips indicating that the NWs are N-polar.

Figure 24 :

 24 Figure 24: 45° view SEM images showing the pc-Si layer (a) before growth of GaN NWs and (b) after NW growth showing deformed pc-Si layer. (c) Top-view SEM image of pc-Si layer with NWs show that the NWs avoid growing on top of the stone-like structures.

Figure 25 :

 25 Figure 25: 45° view SEM image of pc-Si (a) just after crystallisation and removal of Al and (b) after annealing in MBE chamber at the usual NW growth temperature (800 °C) for 3hours, showing the surface is deformed even without Ga and N fluxes.

Figure 26 :

 26 Figure 26: EDS scan for the deformed pc-Si layer shows that there is no presence of foreign material in these stone-like structures and are just made of Si

Figure 27 : 3 . 5 .

 2735 Figure 27: GIXRD 2theta-Chi analysis shows no major change in the fiber-texture quality of the pc-Si layer after deformation.

Figure 28 :

 28 Figure 28: (a) Illustration to show the method to obtain AIC-Si templates for preliminary results of selective area growth (b) top-view SEM image of the pc-Si -SiO2 border after crystallisation and removal of Al. Cross-view SEM images of GaN NWs grown at (c) 800 °C showing NWs are obtained on both SiO2 support and pc-Si layer (d) 810 °C showing NW grow selectively only on pc-Si layer.

Figure 29 : 8 .

 298 Figure 29: Illustration of the steps involved in fabrication of AIC-Si nano-islands by lithography

Figure 30 :Figure 32 :

 3032 Figure 30: Illustration on why it is necessary to use PMMA-MAA combination as photoresist to lift-off AIC-Si layers before crystallisation Figure 31: (a) Top-view SEM image for patterns with numbers ranging from 1-5 (b) 45 o -view SEM image of a single pattern showing the shape of the AIC-Si islands (c) A plot to show the diameter size of each AIC-Si island corresponding to each pattern number

Figure 33 :

 33 Figure 33: 45 o view SEM image of each AIC-Si patterns corresponding to different diameters after the final step of postlift-off cleaning. It can be seen that the patterns are contamination-free with well-defined edges of islands.

Figure 34 :

 34 Figure 34: (a) 45 o view SEM images of patterns of different diameters with NWs on them (b) SEM images showing proof for vertical NW growth on each pattern irrespective of the diameter of the AIC-Si islands (c) Plot to show the verticality and selectivity trend with respect to the diameter of the AIC-Si islands.

Figure 1 :

 1 Figure 1: Bird-view SEM images showing inclined NWs on SiO2 layers, where coalescence can be observed on the base of the NWs. Images are reproduced from Fan et al.7

7

 7 Figure 1: Bird-view SEM images showing inclined NWs on SiO2 layers, where coalescence can be observed on the base of the NWs. Images are reproduced from Fan et al.7

4. 2

 2 State of the art: Silica substrateOn Si (001), the normal growth of self-induced GaN NWs, without any pre-deposition before NW formation, has been shown to result in silicon nitride (SiNx) formation and the authors have observed vertical growth of NWs on it9 . In other studies, where an amorphous layer was deliberately obtained on Si (111) substrate, the growth direction of GaN NWs tends to align with the substrate normal after long growth time but strongly tilted NWs are also observed. The mostly inclined NWs collide rapidly with neighbouring NWs, leading to coalescence and termination of the strongly tilted NW growth. As a result, the angular distribution of growth directions narrows around the substrate normal as growth proceeds. Shadowing effects may also stop the growth of the tilted NWs as the vertical NWs capture most of the incoming vapour species before they can reach the tilted NWs. Fan et al. have grown GaN NWs on 100 nm SiOx layers deposited by plasma enhanced chemical vapour deposition. They obtained inclined NWs on silica where increased growth time reduces the inclination, due to the coalescence of NWs 7 . In another report given by Sobanska et al. for growth of GaN NWs on amorphous Al2O3 layer, through TEM observation it has been shown that NWs coalesce at the very beginning of growth due to high number of nucleation sites on the amorphous layer 8 . However, in all the above-described cases, the tilted coalesced NWs that remain in the bottom part of the NW ensemble produces grain boundaries or dislocations which are detrimental to the optical properties of the sample 10 . First reports of vertical GaN NW growth on non-crystalline substrate was shown by Stoica et al. Highly vertical NWs were grown on ~320 nm thermal silica / Si (100) templates using PAMBE NWs were grown along c-axis 6 .

Figure 2 :

 2 Figure 2: TEM image showing the base of NWs grown on a-Al2O3. Coalescence can be observed by the presence of grain boundaries. Image is taken from Sobanska et al.8

Figure 3 :

 3 Figure 3: TEM images of vertical GaN NWs grown on SiO2 layers. Images are taken from Stoica et al.6

6

 6 Figure 3: TEM images of vertical GaN NWs grown on SiO2 layers. Images are taken from Stoica et al.6

Fig. 5

 5 summarizes the evolution of the NW density with time. As already indicated by the RHEED observations, no NWs were formed for growth times less than 40 minutes. The first NWs are observed for the growth duration of 50 minutes. Then the NW density steadily increases until it starts to saturate after about 3 hours of growth. The longer incubation time on amorphous silica contrasts sharply with the results of Sobanska et al. 11 who compared the nucleation of GaN NWs on crystalline sapphire and on amorphous alumina. These

Figure 4 :

 4 Figure 4: Top row: SEM top views of NWs grown on (a) Si (111) (b) thermal silica and (c) fused silica. The bottom row shows the corresponding cross-sectional views. The scale bar is the same for all panels.

  incubation time) or increases slowly during the 40 minutes of incubation can be answered by considering the evolution of the density of NWs with time. The first hypothesis would imply a constant nucleation probability as soon as the adatom density saturates. Thus, the NW density should increase linearly with time after a short incubation delay.

  observe a long incubation delay followed by a rapid increase of the NW density. This suggests that the net population of resident adatoms increases progressively at the surface (during 40 minutes in the case of silica) until it reaches a critical value. Below the critical density of adatoms, the probability of nucleation is too low. Above this critical density, the onset of nucleation is sudden and all the formed nuclei start to collect adatoms that feed their growth.

Figure 5 :

 5 Figure 5: Variation with growth time of the surface density of the NWs grown on thermal silica substrates, with insets showing the top view SEM image of each sample. The points on the graph are the densities measured for each sample and the solid line is a guide to the eye.

  NWs have random in-plane orientations on the amorphous silica substrates. The ring is narrower for the fused silica, which indicates a better NW verticality as compared to the thermal silica. In contrast, the pole figure of NWs on Si (111) exhibits six well defined spots, which indicate a single in-plane orientation of the GaN NWs in epitaxial relationship with the mono-crystalline Si(111) wafer. The deviation from the perfect vertical alignment is quantified by the cutline analysis of the pole figure (Fig. 6 (b)). A full width at half maximum (FWHM) of 5.2° is obtained for the NWs grown on thermal silica, of 2.1°

Figure 6 :

 6 Figure 6: (a) X-ray pole figures for GaN NWs grown on (i) Si (111) (ii) thermal silica and (iii) fused silica (b) ψ scans (profiles along the ψ axis of the maps in (a)). Dotted lines show Gaussian fits of the experimental values.

Figure 7 :

 7 Figure 7: TEM images of the base region of NWs grown on fused silica showing (a) verticality of the NWs (b) NWsilica interface with wurtzite crystal structure and absence of defects (c) 2D layer between the NWs. (d) Scanning TEM (STEM) image of the NWs grown on fused silica and the EDX mapping of (e) Ga and (f) N elements in NW and 2D layer.

Fig. 7 (

 7 Fig. 7 (a)-(b) do not show any particular feature at the interface, which could indicate that nucleation is

Fig. 8

 8 Fig. 8 shows the atomic force microscopy (AFM) images of the two silica surfaces obtained in tapping mode. The surface features are similar for the two samples. The root mean square roughness values are 0.55 nm and 0.42 nm for thermal and fused silica respectively, so that the angular distribution of NWs ca not be related to the substrate RMS. The thermal silica surface appears slightly wavier (on a scale of hundreds of nanometres) which may explain the broader dispersion of NW directions on this substrate.

Figure 8 :

 8 Figure 8: AFM images of (a) thermal and (b) fused silica.

Figure 9 :

 9 Figure 9: KOH-etched NWs. (a) On fused silica, all the NWs exhibit sharp tips (b) Detached NWs having pointed tips only on one side.

Figure 10 :

 10 Figure 10: SEM images showing the (a) 45° view and (b) top-view of a bush of NWs. (C) EDS data comparison between the bush spot (dark red) and a GaN NW (orange).

  section 2.2.6.4) for IDBs, we observed by TEM dispersed NWs from Si and fused silica to investigate whether the major difference in their PL spectra (namely, the intensity of the 3.45 eV peak) would also correlate with the presence of IDBs. While we observed several NWs grown on Si (111) containing a polarity inversion domain, we did not find this defect in any of the 15 NWs grown on silica that we investigated carefully (Fig.12). Hence, our PL study indicates that, in terms of optical quality, the GaN NWs grown on silica compare favourably with those grown on crystalline Si (111) substrates. This demonstrates that the use of such an amorphous substrate does not result in the incorporation of additional impurities or the formation of nonradiative defects in the body of the NWs. On the contrary, we observe a reduction of the 3.45 eV peak intensity which correlates with a much smaller density (and possibly an absence) of IDBs in the NWs grown on fused silica.

Figure 11 :

 11 Figure 11: Micro-PL spectra from dispersed NWs grown on Si(111) and fused silica.

Figure 12 :

 12 Figure 12: Bright field (a) and dark field (b) STEM images of dispersed NWs grown on fused silica.

Figure 1 :

 1 Figure 1: Illustration showing the (a) covalent bond formation during standard epitaxy of 3D material on 3D substrates and (b) the different combinations of material stacking for van der Waals epitaxy. Inspired from Hong et al. 6

Figure 2 :

 2 Figure 2: GaN NWs grown with graphene as a buffer layer on top of (a) Si (111) (b) Si (100) (c) sapphire and on (d) Silica. Images are taken from Heilmann et al.10,11

Figure 3 :

 3 Figure 3: (a) Top-view and (b) inclined view of GaN films grown on graphene/Si(100) templates using MBE. Images are taken from Araki et al.22

Figure 4 :

 4 Figure 4: Top view SEM images showing (a) small graphene flakes, (b) large graphene flakes on SiO2 templates and (c) a magnified view of the centre of a large graphene flake where several monolayers are stacked.

Fig. 5 (

 5 Fig. 5 (a) shows scanning electron microscopy (SEM) views of GaN NWs grown on a large polycrystalline monolayer of graphene (about 1 cm 2 ). All the NWs are vertical, as seen in the cross-section. Their diameter is around 40 nm. Their in-plane orientation was analyzed by top view SEM observations. Typically, each isolated NW has 6 sidewall facets (m planes) and more complex contours are observed when two or more NWs are merged. Careful observation of the facets reveals that NWs have the same in-plane orientation over large domains of several µm 2 , and this orientation changes randomly from one domain to a neighboring domain. This is illustrated in fig. 5 (b) where different zones of GaN NW in-plane orientation can be observed. Each colored zone contains NWs with the same in-plane orientation and this orientation rotates by a random angle when

Figure 5 :

 5 Figure 5: (a) Cross-sectional SEM image of GaN nanowires grown on a 1 cm 2 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.

Figure 6 :

 6 Figure 6: Top-view SEM image observed at the edge of the ML patch where it can be seen that the NWs grew only on graphene and not on the silica layer (perfect selectivity).

Figure 7 :

 7 Figure 7: (a) Top-view SEM image of NWs on a small graphene flake with perfectly selective growth and (b) magnified version of the same showing in-plane oriented NWs.

Figure 8 :

 8 Figure 8: (a) SEM bird's eye view of the NWs grown near the central part of a pyramidal monocrystalline large 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.

Figure 9 :

 9 Figure 9: SEM images of GaN NWs grown on (a)(c) 50 nm and (b)(d) 300 nm SiO2 carrier layer, taken in (a)(b) 45° titled view and (c)(d) cross sectional view. Panel (a) has the same scale as (b) and panel (c) has the same scale as (d)

Figure 10 :

 10 Figure 10: (a) 2θ/ω scan showing diffraction peaks associated with GaN (0002) and Si (004) planes, indicating that GaN NWs grow along their c-axis. (b) Rocking curve measured for the GaN (0002) reflection. The FWHM (1.62°) characterizes the deviation of the NW orientation from verticality.

Figure 11 :

 11 Figure 11: {112̅ 0} 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.

  Fig. 13b corresponds to a NW cross section with six {101 ̅ 0} facets. Schematics of the GaN lattice orientation is shown in Fig. 13 (d). Blue and red full circles are the N and Ga atoms respectively with their arrangement in the GaN lattice near the edge between two {101 ̅ 0} facets. The combination of Fig. 13 (c) and 13 (d) illustrates the relative in-plane orientation of the two materials: the 〈211 ̅̅̅̅ 0〉 directions of the wurtzite GaN lattice are parallel to the directions of the zigzag chains of

  Figure 12: (a) Photoluminescence spectrum of GaN NWs on a 1 cm 2 patch of 1 ML graphene substrate at 4 K. (b) Gaussian fitting of the near-band edge emission evidencing three spectral contributions at 3.42, 3.45, and 3.472 eV.

Figure 13 :

 13 Figure 13: (a) SEM top-view of a single graphene flake with NWs grown selectively on the flake; inset showing the orientation of the hexagonal flake. (b)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. (c) Schematics of the graphene honeycomb lattice and its zigzag edges at the periphery of the hexagonal flake and (d) schematics of the hexagonal cross-section of a NW with GaN lattice and {101̅ 0} sidewall facet.

Figure 15 :

 15 Figure 15: 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).

Figure 16 :

 16 Figure 16: Morphologies, size, and density of gold nanoparticles on n-layers of graphene after annealing (false-color) (a) Gold nanoparticles on monolayer, bilayer, and trilayer graphene, respectively. (b) Statistics of the size and density of gold nanoparticles on n-layers of graphene.Taken from Zhou et al.40

Figure 17 :

 17 Figure 17: 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.

( 1 ) 2 𝑘

 12 , the nucleation probability then writes: 𝑃 = 𝑃 0 𝑒𝑥𝑝 [ -𝛼𝛾 with 𝑘 𝐵 the Boltzmann constant, 𝑇 the temperature, 𝛼 a geometric constant, 𝛾 the edge energy of the GaN nucleus, 𝛥𝜇 the difference of chemical potential of a (Ga,N) pair between adatom state and solid state in the absence of elastic effects and 𝑃 0 a quantity that varies only slowly with 𝛥𝜇 compared to the exponential term 41 . Elastic effects oppose nucleation by diminishing 𝛥𝜇 of the elastic energy per Ga-N pair (obtained by scaling Equation (1)), with 𝑤 ∞ 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)

Figure 20 :

 20 Figure 20: (a) Top-view SEM image of micro-sized patterned graphene with GaN NWs (b) Zoomed view showing sparse NWs on silica layer (c) Tilted SEM image showing NWs are vertical (d) Bushes found between the graphene patterns (e) Bushes of NWs found on silica layer outside of the patterns.

Figure 21 :

 21 Figure 21: SEM images showing the SAG of GaN NWs on nano-patterned graphene sample.

A. 1

 1 3 provides the RHEED basics for in-situ characterisation. Post-growth characterisation techniques, X-ray diffraction and photoluminescence are briefly introduced in section A.4 and A.5 respectively. Crystal properties of GaN Gallium nitride exists in three crystalline structures: wurtzite, zinc-blende, and rock salt. By epitaxial growth, rock salt form cannot be obtained: GaN crystallizes mainly in zinc blende and wurtzite structures. The zinc-blende structure has a cubic unit cell and the wurtzite structure has a hexagonal unit cell. Due to the cohesive energy per bond difference between wurtzite and zinc-blende (9.88 meV/atom for GaN), wurtzite is more energetically favourable. Wurtzite is also the most commonly observed structure for GaN nanowires grown by plasma-assisted molecular beam epitaxy. Each unit cell contains six atoms of each, Ga and N. The hexagonal unit cell of wurtzite is defined by two lattice constants, a and c, the basal plane lattice parameter and the axial lattice parameter (perpendicular to the basal plane), respectively. The interatomic distance d is described by the internal parameter u=d/c. In an ideal wurtzite structure, the values of the axial to basal ratio and the internal parameter are c/a = 1.633 and u=0.375, respectively. These ideal values can be affected by a small distortion of the tetrahedral angles. Table 1 shows ideal and experimentally observed structural parameters for GaN. The lattice parameters are commonly measured at room temperature by X-ray diffraction (XRD), which happens to be the most accurate one, using the Bragg law. The experimental values can differ from the calculated ones due to various factors like non perfect stoichiometry, stress, free

Figure 22 :

 22 Figure 22: Illustration of a unit cell of wurtzite structure GaN

Figure 23 :

 23 Figure 23: An illustration to show a-, c-, and m-planes in GaN

Figure 24 :A. 4

 244 Figure 24: Illustration showing a line defect having Burger vector (a) perpendicular to the dislocation line and (b) parallel to the dislocation line

Figure 26 :

 26 Figure 26: Illustration to show the RHEED patterns for different type of surfaces

  1

Table 3 -

 3

1: List of metals with their eutectic temperatures having crystallisation temperatures less than 600 °C.

A constant Ar flux of around 30 sccm is supplied to the chamber with a constant pressure of around 10 - 7 mBar.

Acknowledgement

3. Al target is pre-sputtered with Ar to remove the contamination and oxide from the surface of target.

The sputtering is stopped once the voltage of the target stabilises around 380 V indicating the target is pure.

SPUTTERING OF LAYERS

The following steps indicate the default conditions followed in preparing the layers for crystallisation, unless otherwise specified:

1. 30 nm SiO2 / Si(100) support was glued to the holder using a 3M Kapton tape and the substrate holder is kept to rotate at a constant speed to have an homogenous layer. Using DFT calculations and by considering the adsorption and migration energies of various atomic species on these three adsorption sites of graphene, it has been shown that T site is less favorable compared to B-and H-sites 38 .

The in-plane lattice parameters are 0.24682 nm for graphene and 0.3189 nm for GaN. The corresponding 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 (0.98728 nm) and 3 unit cells of GaN (0.9567 nm). 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 fig. 15. 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 (Fig. 15 (a)). Another possibility is obtained by placing one atom on a B site, the other eight being at 5.5 Proof of concept for selective area growth of NWs on graphene

Patterning process of graphene

The observed growth selectivity between the graphene and the SiO2 opens a new way to achieve organized NW arrays. Indeed, by transferring graphene on a SiO2 layer (an amorphous support), which readily hinders the growth of GaN NWs (recall the high growth selectivity that we observe on the sample with graphene flakes), we can pattern the graphene layer to form individual islands and thus obtain selective area growth of NWs on these islands. By adjusting the e-beam exposure parameters, we obtained graphene samples with different motive sizes. In particular, I describe here two samples:

-Micro-size patterned graphene: For this sample the pitch is kept constant (~ 5 µm) and the island size of graphene is changed from ~ 4 µm to ~2 µm.

-Nano-size patterned graphene: For this sample the pitch is varied from 1 µm to 3 µm and the island size of graphene is in the ranges of 30 -300 nm.

A.2 Defects and dislocations in GaN

Due to the hetero-epitaxy, formation of defects in GaN is unavoidable and can be broadly classified into four types based on their dimensionality:

-Point defects: These are zero-dimensional defects associated with a single atomic site like vacancy, interstitial and substitutional atoms.

-Line defects: Defects, like dislocations, associated in one-dimensional manner (with a direction) fall under this category.

-Planar defects: These are defects in two-dimensions associated with a plane or area. Grain boundaries, twins, stacking faults and inversion domain boundaries fall under this category.

-Volume defects: These are three-dimensional defects and are associated to a volume of the crystal. Some common examples are voids, clustering, nanopipes and cracks.

Dislocation causes elastic deformation which leads to the lattice distortion centered around a line which is characterized by the Burgers vector. It described the slip distance in terms of magnitude and distance. Burgers vector can be either parallel or perpendicular to the line of dislocation. Fig. 3 shows an example of both the scenarios. Molecular beam epitaxy allows the oriented growth of thin films on a single crystal substrate, all in a chamber under ultrahigh vacuum. The atoms that arrive at the surface in the form of atomic fluxes, diffuse and interact on the surface to form a thin layer which extends the crystal lattice of the substrate. To do this, highpurity materials are heated from Knudsen cells to temperatures above their melting point, which allows their evaporation in the form of molecular beams and are projected ballistically onto the substrate. The substrate is heated to promote the organization of the crystal. III-and the doping elements are obtained from these solid sources while monatomic nitrogen (N) is generated in a radio frequency plasma cell. The nitrogen source we use is a commercial radio frequency source "unibulb" (PPE or ADDON). The source provides power ranging from 0 to 600 W and the flow rate can be set between 0.2 and 1 sccm (corresponding to a flow rate in cm 3 /min at standard temperature and pressure conditions; 0° C and 1.01 bar), which allows us to modulate the growth rate.

Ideal

The tuning box can stabilize the plasma in the cavity. The growth rate that we use is of the order of 0.27 to 0.30 monolayer per second, which corresponds to a nitrogen flow rate of 0.6 sccm and a power of 300 W.

In this thesis, I used a "Compact12" MBE built by Riber, dedicated to the growth of nitride type semiconductor materials. This growth bed is presented in figure 4. The machine is equipped with gallium, aluminum and indium for III elements source and silicon and magnesium for doping, present in the Knudsen cells and a source of radio frequency nitrogen plasma. To achieve high growth temperatures of GaN nanowires, the machine has a high temperature oven. This allows to grow samples up to 800-820° C for many hours and to carry out temperature flashes up to 1000° C. Temperature measurement is performed by both thermocouple and pyrometer. The latter is placed facing the growth chamber and normal to the surface of the substrate. Growth substrate is mounted on a molybdenum disks called "molybloc" with a hollow center to accommodate the periodicity is visible on the diffraction pattern due to the transmission of the electron beam though these nanostructures. As a result, the diffraction pattern consists of an array of dots.

A.5 X-ray diffraction (XRD)

X-ray diffraction is a powerful method for structural characterization of crystalline materials. In the case of very thin layers, usual diffraction patterns at large angles are not feasible. We use the diffraction pattern of grazing incidence to measure the diffracted intensity from planes perpendicular to the surface. We also acquired pole figures to determine the quality of the crystalline textures.

If we consider a thin isotropic polycrystalline layer, like a powder, it is composed of randomly oriented small crystallites. The direction of the diffracted beam is given by the Bragg law; the angles of diffraction depend on the dimensions of the crystal lattice. Each atomic plane characterized by its interplanar distance is associated to a single diffraction angle for a given length. A typical measurement configuration is the Bragg-Brentano where the x-ray beam incidents on the sample at an incidence angle (θ) relative to the horizontal plane; detector is oriented at the same angle (θ) from the surface and it makes an angle (2θ) with respect to the incident beam.

θ-2θ scanning measures the intensity diffracted by the crystal planes parallel to the surface. This configuration is therefore used to scan the interplanar distances of planes parallel to the surface.

A.5.1 GIXRD

Grazing incidence diffraction can detect planes orthogonal to the substrate, and determine their lattice parameter. The incident beam penetrates the sample surface with a very low incidence angle (0.3 o in our case)

and the diffracted beams are acquired in the sample plane. The emitter-detector plane is almost parallel to the sample surface; thus, only crystal planes orthogonal to the surface may be in the line of observation.

In this configuration, two scan modes are possible:

-2θχ: This method is suitable for the analysis of the planes orthogonal to the surface.

-2θχ/φ: With the additional rotation of φ, parallel planes which are orthogonal to the surface are also detected. This method is suitable for the analysis of polycrystalline thin films.

If the layer is monocrystalline, these two methods are equivalent. However, if the layer has a fibertexture in the plane (that is, the grains which compose it have a prevailing direction parallel to the sample surface), only the 2θχ/φ method will be able to detect it. In the configuration of Bragg-Brentano, it is easy to determine the texture of a thick film if the latter is orthogonal to the substrate surface (that is parallel to the direction of measurement). For example, if the film has a perfect [111] fiber texture, then only its (111) planes diffract. In the case of a thin film observed at grazing incidence, since the measuring direction is parallel to sample plane, the axis of the texture is not observed: only the planes perpendicular to the surface are observed.

If we consider a polycrystalline Si film, the planes that diffract for angles between 2θ = 20 o and 90°, are: {111}, {220}, {311}, {400}, {331}, and {422}. If this film has a fiber texture orthogonal to its surface, only some of these plans families will be able to diffract grazing incidence. For example, for texture [111], only the families {220} and {422} are orthogonal to this axis, the corresponding diffractogram therefore has only two peaks. Unless Another possibility is to fix the sample and tilt the emitter-detector to obtain the angles α and β. The diffractometer that we use in the laboratory allows to use both configurations.

A.6 Photoluminescence

Photoluminescence spectroscopy is a method to measure the light emission due to the recombination of electron-hole pairs induced by laser excitation. Electron-hole pairs are created when the semiconductor is incident with photons of energy higher than the band gap energy (Eg) of the material. After excitation, the following phenomena can occur: In all the cases, the electron-hole pair recombination does not result in photon emission. By the emission of phonons (due to the heating of the crystal), non-radiative recombination is also possible.

The optical properties of the nanowires directly on the as-grown substrate have been probed using the PL setup shown in fig. 8. The PL characterization has been performed on nanowire ensembles placed in a cryostat and excited using a frequency-doubled continuous wave Ar++ ion laser of 244 nm wavelength (λ). Laser spot of approximately 5 μm in diameter was projected by means of a 20X UV microscope objective with a numerical aperture of 0.4. PL spectra were measured using a Jobin Yvon HR460 spectrometer with a 600 grooves/mm grating and a UV-enhanced CCD camera. The optical properties of single nanowire was investigated by microphotoluminescence (μPL) spectroscopy. In order to avoid excitation of multiple nanowires, the nanowires were cut from their substrate by an ultrasonic bath and dispersed with a low density on a Si substrate. The μPL characterization was carried out at 4 K by exciting single nanowires at 244 nm wavelength by a continuous wave frequency-doubled Ar++ ion laser. ABSTRACT -A major advantage of semiconductor nanowires (NWs) is the possibility to integrate these nano-materials on various substrates. This perspective is particularly attractive for III-nitrides, for which there is a lack of an ideal substrate. We examined the use of novel templates for growing GaN NWs by plasma assisted molecular beam epitaxy. We explored three approaches with a common feature: the base support is a cost-efficient amorphous substrate and a thin crystalline material is deposited on the support to promote epitaxial growth of GaN NWs.

In the first approach, we formed polycrystalline Si thin films on amorphous support by a process called aluminum-induced crystallization (AIC-Si). The conditions of this process were optimized to get a strong [111] fiber-texture of the Si film which enabled us to grow vertically oriented GaN NWs. The same idea was implemented with graphene as an ultimately thin crystalline material transferred on SiOx. We illustrated for the first time in literature that GaN NWs and the graphene layer have a single relative in-plane orientation. We propose a plausible epitaxial relationship and demonstrate that the number of graphene layers has a strong impact on GaN nucleation. Proof-of-concept for selective area growth of NWs is provided for these two approaches.

As a simple approach, the possibility of growing NWs directly on amorphous substrates was explored. We use thermal silica and fused silica. Self-induced GaN NWs were formed with a good verticality on both substrates.

Based on our observations, we conclude that the epitaxial growth of GaN NWs on graphene looks particularly promising for the development of flexible devices. Dans la première approche, nous avons formé un film mince de Si poly-cristallin par un procédé appelé cristallisation induite par l'aluminium (AIC-Si). Les conditions ont été optimisées pour obtenir une forte texture de fibre orientée [111] du film de Si qui nous a permis de faire croitre des nanofils de GaN verticaux. La même idée a été mise en oeuvre avec le graphène, matériau cristallin ultimement mince, transféré sur SiOx. Nous avons montré pour la première fois dans la littérature que les nanofils de GaN adoptent une orientation basale bien définie par rapport au graphène. Nous proposons un modèle de relation épitaxiale plausible et démontrons que le nombre de mono-feuillets de graphène a un fort impact sur nucléation des GaN.

Dans ces deux dernières approches, nous démontrons qu'une forte sélectivité de croissance peut être obtenue sur les zones couvertes du film mince cristallin par rapport au support amorphe.

La troisième approche consiste à faire croitre des nanofils directement sur les substrats amorphes. Nous avons utilisé la silice thermique et la silice fondue. Nous avons examiné le temps de latence avant la formation des premiers germes et obtenu des nanofils de GaN de bonne verticalité sur les deux types de silice.

Sur la base de nos observations, nous concluons que la croissance épitaxiale de nanofils de GaN sur graphène est particulièrement prometteuse pour le développement de dispositifs flexibles.

Mots-clés : nitrure de gallium, graphène, épitaxie par jets moléculaires, nanofil, croissance cristalline, cristallisation induite par l'aluminium, substrats de verre, supports amorphes.