Chapter 1

General Introduction

This thesis is a collaboration between the Laboratoire de Photonique et Nanostructures (LPN/CNRS), now Centre de Nanoscience et Nanostructures (C2N), and the molecular beam epitaxy (MBE) manufacturer RIBER. The project belongs to a wider framework, "project E", of the Institut Photovotaique d'Ile de France (IPVF1 ). The later is one of the "Instituts pour la transition énergétique", sponsored by the French government, to nurture and develop partnerships between public institutes (CNRS, local universities) and several industrial French companies (Total, EDF, RIBER, Horiba, . . . ). The IPVF main program advertises the development of solar cells with over 30% efficiency and cost of less than 30 ¢/W in 2030 (i.e the "30-30-30" target). The "project E" is focused on the application and uses of solar cells based on III-V semiconductors in various forms: planar devices, nanowires arrays, tandem solar cells and low cost substrates.

Motivations

The annual energy consumption by human activity was 559.8 EJ (1 exajoule=10 18 J) in 2012, while the potential of the solar energy is 1575-49387 EJ2 . We can thus see that there should be no shortage of energy or energy crisis. However, solar energy is distributed over the whole earth and covers a broad wavelength spectrum.

Many efforts have been made in the past decades to develop efficient solar cells and to reduce the cost of photovoltaic (PV) electricity generation.

The main driving forces for finding low cost, renewable and clean energy sources are the increase of the oil prices, the global warming, the greenhouse effect and the increase of global energy demand [Lewis, 2010]. Initially, the use of sunlight as source of energy was limited to niche applications such as space telecommunications or military applications, but we have now moved towards broad domestic and commercial uses.

The so-called "first generation" solar cells based on thick (100 -200 µm) mono-and multi-crystalline silicon wafers still dominates the current market. Record efficiencies achieved in laboratories are respectively 26.7 ± 0, 5% for mono-crystalline Si and 22.3 ± 0.4 for multi-crystalline Si [Green, 2018]. We also find "second generation" solar cells made of direct bandgap semiconductors such as GaAs, CuIn x Ga 1-x Se 2 or CdTe, but they have yet to replace their silicon counterparts.

GaAs is a good candidate to reach higher conversion efficiencies than silicon. The later has an indirect energy gap of 1.12 eV (which corresponds to a wavelength of 1107 nm) while GaAs has direct bandgap of 1.424 eV, (871 nm). As a consequence, a Si cell requires about 100-200 µm thick absorber, while a GaAs need only 2 to 4 µm to completely absorb the maximum number of photon. Green et al. reports that the record efficiency for a GaAs single solar cell is of 28,8% [Green, 2018], which is the most efficient single-junction solar cell developed to date. This particular cell is composed by an absorber about 2 µm-thick, epitaxially grown on a GaAs (100) substrate and then transferred to a host flexible substrate [START_REF] Kayes | 27.6% Conversion efficiency, a new record for single-junction solar cells under 1 sun illumination[END_REF].

To compensate for the tenfold price difference between Si and GaAs substrates, the main idea is to recycle the expensive GaAs substrate so it can be re-used for a large number of epitaxial growth. Without the substrate cost, the GaAs-based technologies are expected to be more competitive with respect to the traditional Sibased technologies. Despite the achievement of record efficiency on single cell, the substrate recycling technology is still limited and it has yet to prove commercially viable. Alternatively, a simpler approach to reduce the substrate cost is to simply eliminate the starting mono-crystalline wafer.

Proposal

In this work, we investigate an alternative way to fabricate low-cost substrates for the epitaxial growth of GaAs thin-films. These substrates are composed of a uniformly oriented poly-crystalline Ge layer made of thin Ge crystalline patches. These pseudo-substrates are fabricated by the Metal Induced Crystallization (MIC) of an amorphous germanium layer deposited on a glass substrate. The MIC process allows for an oriented crystallization of the amorphous Ge at relatively low temperature (250 -350 • C), promoted by the presence of a metal. We choose Ge because it is near lattice matched to GaAs. Nevertheless, the MIC process also works with other materials such as silicon.

Many metals can be used to promote MIC but the particular case of Aluminium induced crystallization (AIC) has attracted much interest. The low cost of Al and the low annealing temperatures required for AIC of amorphous Ge are very attractive. In the literature, Si and Ge crystalline layers obtained from AIC have been investigated as potential candidates for various devices, including solar cell absorbers [Gordon, 2007;Gall, 2009], transistors [Chen, 2015] or buffer layers for epitaxy [Toko, 2015;Cohin, 2013].

Thesis structure

The final goal of this work is to realize a solar cell on a glass support using a MIC-Ge layer as a pseudo-substrate. This work is divided into three main chapters which correspond to different challenges (Figure 1.1). • In chapter 2, we first investigate the process of MIC of amorphous Ge. In particular, we develop a new in situ optical microscope to follow the kinetics of the crystallization processes.

MIC poly-Ge

• In chapter 3, we investigate the epitaxy of GaAs layers by MBE on Ge wafers and on MIC-Ge pseudo-substrates. We characterize twin defects specific to the (111) orientation of the substrates.

• In chapter 4, we fabricate and characterize GaAs-based solar cell devices on different substrates from monocrystalline GaAs (111)B and Ge(111) wafers to MIC-Ge pseudo-substrates.

In each chapter, we summarize the relevant state of the art before describing our experimental findings.

Chapter 2

The low-cost substrate using Metal Induced Crystallization

State of the art

We first review the various techniques that can be used to crystallize amorphous Germanium (a-Ge).

Typical crystallization techniques

Among the possible techniques which can be used for the crystallization of a-Si or a-Ge we find: Solid Phase Crystallization (SPC) [Yamaguchi, 2001;Tsao, 2009;Takahara, 2018], laser annealing [START_REF] Im | Phase transformation mechanisms involved in excimer laser crystallization of amorphous silicon films[END_REF]Eisele, 2003;Watakabe, 2004;Kühnapfel, 2015] and Metal Induced Crystallization (MIC) [Wang, 2015;Katsuki, 2001]. Typical SPC crystallization temperature is about 450-600 °C [Gaudet, 2006;Tsao, 2009]. SPC of Ge on amorphous substrates results in small grains (tens of nm) [Tsao, 2009]. MIC allows one to obtain polycrystalline semiconductor thin films with larger grains (tens or hundred of µm) [Nakazawa, 2013b;Higashi, 2015] at lower temperature (125-350 °C for Ge with Al), always below the eutectic temperature between the semiconductor and the metal. Also, MIC yields a crystalline texture, (111)-oriented for thin films [Toko, 2012], while the SPC usually produces randomly oriented grains [Yoshimine, 2018]. Laser annealing allows for local heating and can produce large grains with specific crystal texture but requires special equipment.

Solid Phase Crystallization

With SPC, it is possible to crystallize amorphous Ge deposited on an amorphous substrate such as glass. Tsao et al. showed that the transition temperature of Ge (300 nm, sputtered on glass at room temperature) is between 400 °C and 500 °C for a 20 h anneal [Tsao, 2009]. Better crystallinity is obtained however with SPC at 600 °C.

The mechanism of underlying SPC is that the breaking of bond is thermally activated at defects sites in the amorphous material. This results in an activation energy for the SPC of amorphous Ge of 2.0 eV [Yamaguchi, 2001], which is very high compared to the typical annealing temperature (typically 450-600 °C) but still permits measurable crystallization rate. Toko et al. deposited an amorphous Ge layer (50 nm) in a MBE machine and then annealed the Ge at 425-500 °C for SPC [Toko, 2009]. On the one hand, these authors found that the resulting Ge is actually p-type self-doped and that the hole concentration decreases with the annealing duration and temperature. On the other hand, the hole mobility also decreases with the annealing temperature but it increases with the annealing duration. Larger grains (∼ 200 nm) were observed with a SPC temperature of 425 °C as compared to 500 °C (∼80 nm grain size). These large grains were correlated with a higher hole mobility. To maximize the hole mobility and decrease the defect density, these authors found that the best results are obtained with a two-step annealing: a first step at 425°C to crystallize large grains and a second step at 500 °C to decrease the defect density. They get a hole mobility of 140 cm 2 /(V • s) and a hole concentration of 5 • 10 17 cm -3 .

To further increase the hole mobility, Toko et al. [Toko, 2017] combined the SPC with a preliminary step of heating the substrate during the deposition. Below 150 °C the Ge does not nucleate during the deposition. Still the atomic density of the a-Ge increases with the substrate temperature below this value. The combination of a deposition temperature of 125 °C and a very low SPC temperature of 375 °C (for 140 h) allows them to obtain grains as large as 5 µm and hole mobility of 340 cm 2 •V -1 •s -1 . Experiments also reveal that the hole mobility is more dependent on the deposition temperature than the grain size.

Further improvements of the SPC of Ge were achieved by Yoshimine et al. [Yoshimine, 2018]. They found that the Ge thickness (50-500 nm) deposited from a Knudsen cell (MBE) affects the final quality of the crystallized Ge. The best result was obtained with a Ge thickness of 400 nm and by combining the two temperature annealing [Toko, 2009] (first annealing at 450 °C 5 h and post-annealing at 500 °C for 5 h) with the heating of the substrate during the deposition [Toko, 2017] (150 °C). In this way, these authors obtained a hole mobility of 450 cm 2 /(V • s). However, the obtained crystalline grains show no crystal texture according to electron backscattered diffraction.

Laser annealing

Laser crystallization of Ge or Si on amorphous substrates can take different names depending on the technique: laser assisted liquid phase crystallization (LPC) [Kühnapfel, 2015], excimer laser crystallization (ELC) [Liao, 2017], excimer laser annealing (ELA) [START_REF] Yeh | Superlateral growth of a-Ge film by excimer laser annealing[END_REF], semiconductor diode laser annealing (SDL) [START_REF] Sakaike | Crystallization of amorphous Ge films induced by semiconductor diode laser annealing[END_REF], pulsed laser annealing [START_REF] Watakabe | Electrical properties for poly-Ge films fabricated by pulsed laser annealing[END_REF] and laser interference crystallization (LIC) [Eisele, 2003].

Eisele et al. [Eisele, 2003] used a pulsed Nd:YAG laser and found that an homogeneous laser profile gives a poor crystalline structure with 50 nm grain size. Instead, a three beam interference pattern with a period of 6 µm yields a narrower molten line and a larger lateral intensity gradient. This allows them to crystallize large grains up to 2 µm.

Sakaike et al. [START_REF] Sakaike | Crystallization of amorphous Ge films induced by semiconductor diode laser annealing[END_REF] worked with a continuous wave 807 nm laser. These authors were able to control the crystallization by changing the laser line speed from 400 nm/s to 100 nm/s. At a speed lower than 170 nm/s they observe a melting and recrystallization, while for speeds higher than 200 nm/s they observe SPC. The crystal grows laterally in the <111> direction to minimize the solid-liquid interface energy. This results in a [220] preferential orientation in the vertical direction. They obtained grain size of 200 nm by 1 µm Liao et al. [Liao, 2017] were able to increase the grain size to 1 µm by changing the power of the excimer laser. Using a similar excimer laser (XeF pulse width 25 nm) Yeh et al. [START_REF] Yeh | Superlateral growth of a-Ge film by excimer laser annealing[END_REF] increased the grain size from 0.5 µm to 4.1 µm by changing the absorption coefficient of the silica underlayer. This was possible by depositing SiO x N y and SiO 2 layers on the glass substrate. Similarly, Watakabe et al. used a XeCl excimer laser. They obtained 126 nm grain size. Kuhnapfel et al. [Kühnapfel, 2015] show that for Si crystallization with a continuous wave laser, the crystalline texture can be controlled with a buried seed layer. In the scanning direction, the grain size reached the sample size of 5 cm.

Metal Induced crystallization

Metal Induced Crystallization (MIC) is the crystallization of an amorphous material in contact with a metal. The annealing temperature is always below the eutectic temperature of the metal-semiconductor mixture, so that the layer remains in solid phase. The main differences with SPC are that the metal acts as a catalyst for the crystallization and that the annealing temperature is typically lower.

Metal Induced Crystallization of Silicon

The MIC technique allows one to obtain thin polycrystalline films of Si or Ge from amorphous layers. The amorphous semiconductor layer and the metal can be easily deposited by standard evaporation or sputtering techniques.

Inversion mechanism

The reaction of amorphous Si or Ge with a metal is known from the early 1970s [Oki, 1969]. However, only in 2000 did Nast et al. [Nast, 2000b] give an explanation of the mechanism for the crystallization of an a-Si/Al layer deposited on an amorphous substrate (glass). Following the work of Nast et al. , this Aluminium induced crystallization (AIC) is also referred to as the ALILE process, for Aluminum-Induced Layer Exchange, because the Al and amorphous Si layers exchange positions grain by grain as the Si crystallizes. Different steps could be identified in the annealing (at 375-525 °C for Si in the experiment of Nast et al. ; (left). Cross sectional SEM micrograph at these successive MIC steps (extracted from [Nast, 2000b]) (right) inside the metal or at the lower Si interface reaches a critical value [START_REF] Sarikov | A kinetic simulation study of the mechanisms of aluminum induced layer exchange process[END_REF].

3. Then, the crystalline grains grow laterally inside the Al layer, due to continuing diffusion of Si into Al during the annealing process (Figure 2.2 (c)). Meanwhile, the growing Si crystal is constrained between the Al grains and the substrate. This creates stresses and eventually causes the outward diffusion of Al from its original position, into the a-Si layer.

4. If the a-Si is initially at the top and the Al at the bottom, the crystalline Si (c-Si) grows at the bottom and the Al moves to the top (Figure 2.2 (d)). The inversion mechanism also works if the a-Si is initially below the Al; the final c-Si is then above the metal [Nast, 2000a].

From the thermodynamic point of view, the system free energy decreases due to two phenomena: the crystallization of the amorphous Si material and the change of interface energies.

When the semiconductor is in contact with the metal, the strong covalent bond Si-Si (or Ge-Ge) is weakened due to an electronic screening, as reported by A. Hiraki et al. (see [Wang, 2015]). The interaction can extend up to two monolayers and causes an intermixing of the two materials at the interface. It typically follows that the Si (Ge) atoms can easily diffuse and wet the Al grain boundaries, replacing the initial metal-metal interface with two metal-semiconductor interfaces [Wang, 2007;Wang, 2008b;[START_REF] Sarikov | A kinetic simulation study of the mechanisms of aluminum induced layer exchange process[END_REF].

When a high concentration of Si (or Ge) atoms inside the Al is reached, the crystal nucleation can occur at three locations: at the Al grain boundaries, at the metalamorphous semiconductor interface or at the metal-substrate interface. Depending on the interface surface energies or the nucleation site, different behaviors (crystal orientation, kinetics) can be observed [Wang, 2008b;Wang, 2006].

Regular arrays of crystalline patches

Recently, Cohin et al. [START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF] have proven that the MIC-Si can be used to obtain thin monocrystalline patches with defined sizes, shapes and positions. However, the size of the mono-crystal patches cannot exceed the size of the Si grains of a continuous layer (1 µm in their case); above this, polycrystalline Si patches are formed.

The patches are highly (111) oriented (with a misorientation of only 7%) even if there is sometime a parasitic Si double layer. The maximum size of fully monocrystalline patches depends also on the annealing temperature. These authors report a threshold of 450 nm at 500 °C, 700 nm at 350 °C and 1 µm at 250-300 °C. This is consistent with the size of grains obtained by the MIC of Al/Si continuous layers, as measured by the variation of the density of grains with temperature.

The same authors found that the radial growth velocity and the nucleation rate have similar activation energies of 0.9 eV and 1.1 eV respectively, and concluded that the kinetics are not affected by the size of the patches.

Nonetheless, these authors found that smaller patterns have lower probability to nucleate and crystallize at 250 °C, below the expected decrease due to the reduction of the surface area. They attributed this effect to a partial damaging of Al during the fabrication process. The normal rate can be restored by increasing the annealing temperature to 300-350 °C, which they attributed to a recrystallization of the Al layer.

Finally, these authors were able to obtain 10-20 nm thick squares pads of (111)oriented crystalline Si, with a lateral size between 170 nm and 7.5 µm. (c-e) Gaussian model for patch amorphization during the etching of the trenches between the patches (extracted from [START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF]).

Metal Induced Lateral Crystallization

Si 1-x Ge x alloys with 0<x<1 can crystallize by a Metal Induced Lateral crystallization (MILC) process in the presence of nickel (Ni). For example, Ni patterns (5 nm thick) deposited (by optical lithography) on 100 nm thick amorphous (Si,Ge) [Kanno, 2006] yield flat crystalline grains wider than 10 µm at an annealing temperature of 600 °C, which is 50 °C below the spontaneous nucleation temperature (SPC nucleation). c). (extracted from [Kanno, 2006]). crystallization above and below the spontaneous nucleation temperature. Below this temperature, we see the polycrystal growth around the Ni pattern (Figure 2.4 (b)). The crystal also expands with annealing time at constant velocity. Above the spontaneous nucleation temperature, no lateral growth is observed because SPC is faster than the MILC growth. Figure 2.4 (c) shows the temperatures at which the MILC is possible as a function of the Ge content (x) of the alloy. Oki et al. were the first to report the lowering of the crystallization temperature of amorphous germanium in contact with a metal [Oki, 1969]. However, the detailed mechanism of MIC-Ge has only been described in the last ten years, especially by Toko and coworkers.

Metal Induced Crystallization of Germanium

Metals candidates for the MIC of a-Ge

Knaepen et al. [Knaepen, 2009] have investigated about twenty transition metals for the MIC of a-Ge films.

Figure 2.5 -Crystallization temperature as function of the melting temperature of a selection of eutectics and germanides formed between various semiconductors and metals [Knaepen, 2009]).

Three of these transition metals (Au, Al and Ag) form an eutectic compound with Ge, instead of solid germanide phases. The systems with lowest eutectic temperatures are Ge/Al (420 °C, [START_REF] Mcalister | [END_REF]) and Ge/Au (361°C [Okamoto, 1984]). The corresponding phase diagrams are shown in Figure 2.6 (a) and (b) respectively.

Typically, the lower the eutectic temperature, the lower the crystallization temperature for the MIC. This relates to the metal-assisted weakening of Ge covalent bonds [Knaepen, 2009] which is beneficial for both processes. The lowest MIC temperature of a given metal/Ge couple is typically equal to two thirds of that of the lowest temperature germanide or metal/Ge eutectic (Figure 2.5). [START_REF] Mcalister | [END_REF]); (b) Phase diagram of the Au-Ge system (from [Okamoto, 1984])

We choose aluminium (Al) for our experiments rather than gold (Au) for the following reasons:

• availability and lower cost of Al

• simple way to create on oxide interlayer (AlO x ), see section new section 2.1.3.2.

• large body of existing data in the literature

The interlayer

The diffusion of Ge into the metal can be controlled by the insertion of an "interlayer" between the Al and Ge materials, which acts as a diffusion barrier [Kurosawa, 2009a]. This layer is typically a metal or a semiconductor oxide. For Al-MIC, the obvious choice is AlO x , created by the oxidation of the top surface of the Al layer. The addition of the interlayer usually results in the preferential orientation, i.e. crystal texture, of the Ge grains, due the nucleation at the Al/SiO 2 rather than inside the metal grain boundaries [Kurosawa, 2009a;Toko, 2014b], see Figure 2.7.

Ge can also form an oxide which can be used as an interlayer to control the diffusion of the semiconductor in the metal. This was investigated by Numata et al. [Numata, 2014b] and Hu et al. [Hu, 2010].

In the case of the Al-MIC of amorphous Ge, different behaviour were observed if the crystallization proceed with or without an interlayer. Without interlayer, the outcome is a mix of the two initial layers with a random crystallographic texture [Hu, 2010].

Detailed analysis of crystallized Ge layers without an interlayer revealed that they consist of small randomly oriented grains, due to nucleation inside the metal (away Right: schematics of the expected Ge concentration inside the layers (extracted from [Toko, 2014b]) from the substrate and the interlayer interfaces) [Park, 2012]. The presence of the interlayer suppresses nucleation in the metal grain boundaries [Kurosawa, 2012]. Instead, if the Al layer is thin enough (less than 50 nm), the nucleation occurs preferentially at the interface between the metal and the substrate, which typically leads to the preferential (111) orientation of the crystalline Ge layer.

The first experiments with Al [Kurosawa, 2009a] demonstrated the need of the oxide interlayer in order to obtain large oriented grains. Kurosawa et al. optimized the oxidation time of the Al layer to form the AlO x interlayer [Kurosawa, 2009a]. In combination with layer thickness optimization, they experimented oxidation duration between 5 min and 24 h [Kurosawa, 2012]. 24 h oxidation creates a thick oxide interlayer, which delays the layer-exchange process.

The combination of medium oxidation time (30 min) and low temperature (325 °C) gave the best (111) crystalline orientation (99 % of the grains) (Figure 2.8). Ge grains can also show high orientation fraction with an oxidation time of 5 min and the same annealing temperature (325 °C) [Toko, 2012]. Later, Numata, Toko et al. generally used oxidation times of 5 or 10 min.

Physically, the AlO x interlayer reduces the diffusion rate of Ge atoms in metallic Al by 8 orders of magnitude at 325 °C (from 1.19 10 11 cm 2 •s -1 to 2.09 10 19 cm 2 •s -1 ) [Toko, 2014b]. Decreasing the diffusion rate of the semiconductor increases the time required to complete the inversion so that the introduction of an interlayer requires an adjustment of the annealing time.

The use of a thin AlO x , with 2 nm of Al deposited on the top of Ge layers, is also possible in the case of "reverse MIC" [Nakazawa, 2014a]. Here also, the grain size increases with the increasing oxidation time of the AlO x interlayer, leading to a fraction of 93% of (111)-oriented Ge grains, up to 46 µm in diameter, for 180 min of exposition to air.

Concerning the GeO x interlayer, longer oxidation duration in air is required with [Toko, 2012]) respect to AlO x . Oxidation of 1 nm of Ge on top of Al was carried out for 24 h by Numata, Toko et al. [Numata, 2014b;Toko, 2014a]. In contrast, oxidation of germanium interlayer with ozone require less than 3 min [Hu, 2010] or even less than 1 min [Hu, 2012;Hu, 2015].

, (c) 100 h, (d) 10 h, (e) 30 h, (f) 100 h, (g) 30 h, (h) 100 h, and (i) 400 h. (extracted from

For Au-MIC, no attempt has been made to create a GeO x interlayer. An additional aluminum oxide interlayer was found to be effective in controlling the diffusion of germanium in gold [Park, 2013]. Park et al. proved that by tuning the thickness of the aluminum oxide from 0 to 10 nm, it is possible to optimize grain orientation and size [Park, 2013]. For AlO x 0 to 6 nm thick, small grains with random orientation appeared. For 7 nm thick AlO x , large (111)-oriented grains (> 50 µm) were obtained by Park et al. (2013) [Park, 2013]. For AlO x thicker than 8 nm, the inversion process is inhibited. Further optimization was performed by Higashi et al. [Higashi, 2015] by depositing a 0.6 nm Al 2 O 3 by Atomic Layer Deposition (ALD) instead of sputtering. These authors obtained grains larger than 500 µm.

The use of AlOx as an interlayer

• leads to a good (111) texture of the Ge layer;

• affects the kinetics of crystallization;

• requires specific optimization if performed by direct oxidation of the Al layer.

The best results are obtained for an AlO x interlayer associated with long annealing times and low annealing temperatures. Germanium oxide (GeO x ) can also be used as an interlayer.

The temperature and duration of (MIC) annealing

The annealing temperature is a key parameter in the MIC process. The main difference between Al-MIC and Au-MIC is that the first one requires a higher temperature, due to the respective difference in eutectic temperature between Ge and the two metals [Knaepen, 2009].

Increasing the annealing temperature increases the diffusion of the semiconductor inside the metal and the overall crystal growth rate. However, it also degrades the crystal texture and the well-defined (111) crystal orientation is progressively lost [Toko, 2012;Nakazawa, 2014a;Toko, 2013b]. This dependence can be explained as follows: higher temperature corresponds to a lower saturation threshold for the Ge nucleation inside the metal grain boundaries [START_REF] Sarikov | A kinetic simulation study of the mechanisms of aluminum induced layer exchange process[END_REF]. Consequently, randomly oriented Ge nuclei appear because the Ge atoms crystallize before reaching the lower metal-substrate interface where the oriented crystallization occurs [Toko, 2013b].

Toko et al. [Toko, 2013b] investigated the influence of the temperature on the Alinduced crystallization of Ge, in the range of 325 -400 °C. These authors show that lower temperatures result in a high orientation fraction, whereas, high temperature annealing leads to a completely random texture. As best, they obtained 98% (111) oriented Ge grains with an average diameter of 30 µm, at 325 °C. No further decrease of temperature was suitable in this case, because the annealing duration was already 100 h. Numata et al. [Numata, 2014a], attempted to decrease the annealing temperature by changing the sample structure. They introduced a new layer (1-10 nm thick Ge film) between SiO 2 and the Al layer. This underlayer promotes Ge nucleation due to a faster supersaturation of Ge at the Al/substrate interface. This results in a large density of grains but a slower lateral growth and finally smaller grains. The best compromise is to use no more than 1 nm amorphous Ge as underlayer. By combining this thin underlayer with a lower annealing temperature, these authors obtained large grains (above 100 µm) at 275 °C and a high (111) orientation (99 % fraction).

Alternatively, Numata et al. also succeed in decrease the annealing temperature down to 200 °C by substituting the AlO x interlayer with the GeO x interlayer [Numata, 2014b]. Toko et al. [Toko, 2014a] managed to grow oriented Ge grains at 180 °C by combining a Ge underlayer of 1 nm between the SiO 2 and the Al layer and a GeO x interlayer .

The temperature and annealing duration required to activate the MIC process heavily depend on the sample structure (layer thickness, chemical nature of the interlayer, type of metal . . . ). Specific optimizations are required for each variant of MIC structure. Kurosawa et al. and Toko et al. studied the optimization of the thickness of the stacked Al/Ge [Kurosawa, 2012;Toko, 2013a] or Al/Si [Kurosawa, 2009b;Kurosawa, 2014].

The thickness of aluminium and germanium layers

These authors observe that reducing the metal layer thickness below 100 nm causes the suppression of the random bulk nucleation and the apparition of the crystal texture in the Ge or Si layer. For thicker Al layers, the substrate has no influence on the Ge crystal, which grains are always randomly oriented. High (111) orientation can be obtained by thinning the Al layer down to 50 nm [Kurosawa, 2012]. (50, 75, 100 nm). The (Si,Ge) layer has the same thickness as the Al layer (extracted from [Kurosawa, 2012]) Nakazawa et al. [Nakazawa, 2013b] investigated how the Al/Ge thickness ratio influences the orientation of the final crystallized Ge. They found that a ratio > 1 results in oriented grains but with a poor coverage. Al/Ge ratios < 1 led to an improved surface coverage but with randomly oriented grains. The ratio 1:1 is thus a good compromise. These authors also report that the Al layer should be thicker than 30 nm to avoid the inhibition of the inversion process.

Combined with the results of Toko et al. [Toko, 2013a], these observations limit the Al thickness to between 30 nm and 100 nm. The best compromise in terms of coverage and orientation fraction is to use a 50 nm thick Al layer, coupled to an AlO x interlayer (5 min oxidation in air) and an annealing temperature of 350 °C. In these conditions, Nakazawa et al. report a fraction up to 97% of (111)-oriented grains and an average grain size of 70 µm [Nakazawa, 2013a].

As regards the Au-MIC of Ge, Park, Sadoh and coworkers study a-Ge/Au stack of 100 nm [Park, 2013;Sadoh, 2015;Sadoh, 2016] or 50 nm [Park, 2014a;Park, 2014b] and also succeed in obtaining (111)-oriented Ge grains, as large as 50-100 µm.

• The thickness of the Al and Ge layers must be similar.

• Good crystallographic texture is obtained for individual Ge and Al layers thinner than 50 nm.

• Very thin layers (below 30 nm) may be difficult to crystallize.

The type of substrate

For Ge-AIC, the main substrates used are quartz or fused silica [Toko, 2012;Peng, 2012], but the use of thermally grown silica on Si (100) is also possible [Hu, 2012].

Others, such as Park et al. [Park, 2014a], focus on plastic (polyimide) substrates to take advantage of the lower annealing temperature.

In particular, Oya et al. report a fraction of 95% (111)-oriented Ge grains with an average size of 100 µm when using polyimide coated with a 100 nm SiO 2 layer, using a 45 nm thick Ge layer, an AlO x interlayer and an annealing temperature of 325 °C [Oya, 2015]. In this case, a relation between the thickness of the plastic substrate and the grain orientation appeared. If the plastic substrate it too thin, the annealing results in the bending of the substrate, which facilitates Ge nucleation in the metal layer and the nucleation of randomly oriented Ge crystals. However, a SiN layer can be used to strengthen the plastic substrate and to recover a high (111) orientation fraction of the grains [START_REF] Oya | Direct synthesis of highly textured Ge on flexible polyimide films by metal-induced crystallization[END_REF].

In the more general case, if the substrate or the "supporting layer" is amorphous, the crystallization usually yields (111)-oriented nuclei because it is the most energetically favorable orientation [Stekolnikov, 2005]. However, experiments using Au-MIC show that alumina-covered (Al 2 O 3 ) quartz substrates can yield a (100)-oriented Ge layer [Toko, 2014b].

It is possible also to crystallize Ge on a "conductive supporting layer", which can latter play an active role in the final device. Layers of materials such as TiN [Toko, 2014b;Nakazawa, 2014b], AZO (Al-doped-zinc-oxide) and ITO (indium-titaniumoxide) [Nakazawa, 2014b] have been studied for solar cell applications to act both as the back contact of the cell and as the MIC supporting layer. It is important that these materials are thermally stable and do not diffuse into the Ge crystals. The best results are obtained with a TiN supporting layer (70 nm) on a glass substrate, which yield (100 µm) wide grains and high (111) orientation (96%).

For Au-MIC, Park,Sadoh et al. succeeded in crystallizing (111)-oriented Ge directly on plastic substrates with softening temperature of 300°C [Sadoh, 2015;Park, 2014a;Park, 2014b]. They obtained large grains about 50 µm wide [Park, 2014a]. The results were comparable when either using polyimide wafers coated with silica or pure silica wafer.

• Various substrates can be used as long as they do not react with the metal or germanium, the most common being SiO 2 .

• Plastic substrates can be used provided the annealing temperature is kept low enough.

• The chemical nature of the substrate can affect the crystallographic texture of the crystallized layer.

Conclusion

The Metal Induced Crystallization technique (MIC) is a promising method to obtain large-grain (50 µm) polycrystalline Ge films on low cost substrates.

Possible substrates explored in the literature are SiO 2 , quartz, TiN, AZO, Al 2 O 3 , flexible plastics (mainly polyimide). All these substrates could allow a cost reduction in the fabrication of devices, especially as replacement of III-V substrates.

While there is a vast choice of possible metals for the MIC of amorphous Ge, the most promising and the most investigated are aluminum and gold combined to a silica substrate.

In addition to the main Al (or Au) and Ge layers, it is necessary to introduce an interlayer (AlO x ) and an underlayer (Ge) to control the crystal texture and the kinetics of the MIC crystallization, respectively.

Many authors use Al-MIC because an AlO x interlayer can be formed simply by exposing the metal layer to air. In some experiments, the AlO x interlayer is replaced by a GeO x interlayer.

In the literature, flat, crystallized and (111) oriented germanium platelets more than 50 µm wide have been obtained, using Al-MIC of amorphous Ge, through careful tuning of the interlayer and optimization of the annealing conditions. These platelets may be suitable candidates for the subsequent epitaxy of GaAs and solar cell fabrication, which will be explored in the next chapters.

Sample fabrication

In this section, we present the workflow used to fabricate the MIC samples. We first obtain a silica surface, either from the dry oxidation of silicon wafers, or by the plasma enhanced deposition of silica, or directly from a bulk silica wafer. Then we depose and process the various layers (Al, Ge, interlayer, etc. . . ) using e-beam evaporation. Finally, we anneal the sample to induce the crystallization of Ge and we etch the remaining metal to obtain pure Ge crystalline platelets.

Overview

The fabrication steps of the samples are shown in Figure 2.10. All materials are deposited in the same e-beam evaporator, which allows us to switch material sources without exposing the sample to atmosphere. Starting from a clean silica surface, we first deposit a 0-2 nm Ge underlayer, followed by 15-40 nm Al which is coated with a 0-1.5 nm thin Ge layer.

The samples are then exposed to air to form an AlO x or AlGeO x interlayer, before the deposition of the final 16-40 nm Ge layer (at a rate of 1 nm/s). The samples are then annealed in a trizone oven or in a closed programmable hotplate under nitrogen flow. After the crystallization, the Al is removed by wet-etching. (c-Ge). A thin (1 nm) amorphous Ge layer is used as an underlayer. The AlO x or GeO x interlayer is obtained by air oxidation of the Al layer (from [START_REF] Pelati | In Situ Optical Monitoring of New Pathways in the Metal-Induced Crystallization of Amorphous Ge[END_REF]).

Type of silica

Toko et al. mainly use quartz or fused silica as MIC substrates. We use a wider range of silica: thermal silica, fused silica and silica grown by plasma enhanced chemical vapor deposition (PECVD). In principle, the MIC process should only be impacted by the top surface of the silica.

In this work, we mainly use thermal silica obtained by dry oxidation of Si (100) wafers. We perform the oxidation in a fused quartz tube oven for 15 min at 1050 °C under an O 2 flux (dry oxidation). The oxide grows at the interface of the Si substrate and already formed silica via the diffusion of oxygen molecules inside the silica layer. The typical oxide thickness is 40-60 nm, as measured with an optical reflectometer.

Alternatively, we sometimes use bulk silica 2" wafers (JGS2 grade, from Univer-sityWafer Inc). These are double side polished (DSP).

All silica surfaces from bulk wafers are HF-treated (exposure to HF at 5% for 3 min, followed by rinsing for 2 min in deionized water) prior to Al and Ge deposition in order to remove possible surface contamination.

We also used silica grown by PECVD on a Si (100) wafer. The radio frequency (RF) PECVD reactor is a Unaxis D200. The deposition temperature is 280 °C.

Compared to bulk silica substrates, the samples fabricated using silica on Si (100) wafer are easier to cleave and to characterize. They also allow for a more precise temperature measurement with a pyrometer during subsequent epitaxial growth in the molecular beam epitaxy (MBE) chamber. The silica substrates are fully transparent and can be used to characterize the MIC process from both sides of the bilayer during crystallization (see section 2.4.2.1). An advantage of the fused silica wafers over common glass is that they can be heated in the MBE machine at 600-640 °C.

Deposition of the semiconductor and metal materials

Common deposition methods for a-Ge or a-Si are magnetron sputtering [START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF]Toko, 2012;Nast, 2000b] and evaporation [Park, 2012;Wang, 2009;Kurosawa, 2012]. We use the same e-beam evaporation machine (PLASSYS MEB 550 SL) to deposit both Al and a-Ge (Figure 2.12)

The e-beam evaporation machine has a rotating housing that can host different crucible liners containing different materials such as gold, aluminum and germanium. The rotating housing allows us to switch from one material to another without breaking the vacuum in the chamber. The material to be deposited is heated by electrons generated by a thermionic filament. Electrons extracted from the heated filament are directed to the crucible by a magnetic field, thus causing the melting and the evaporation of the material to be deposited. The crucibles are water-cooled to limit degassing and contamination of the hosted element by the crucible material. The evaporated material condensates on the sample mounted upside-down and positioned on the crucible. Before and after the evaporation, a shutter protects the sample from contamination. The deposition rate is monitored with a quartz balance. Aluminium is typically deposited at a rate of 1-2 nm/s and Ge at a rate of 0.01-0.1 nm/s for thin layers (underlayers or interlayer) and 1 nm/s for the thick top layer.

WATER-COOLED HEARTH

Low temperature annealing

We use two types of ovens to anneal the samples. One is a trizone fused quartz tube and the other is a closed hotplate. In both cases, heating is performed under a nitrogen flow. The closed hot plate will be described in detail in section 2.4.1.1. The trizone oven (from AET) is designed to provide a homogeneous temperature along the fused quartz tube by regulating the temperature using three independent heating systems and thermocouples. In such way, the oven can compensate the losses of heat from the two extremities. For each heating sub-system, the temperature is controlled by a proportional-integral-derivative (PID) controller which compares the set and measured temperatures and adjusts the power heating. Schematics of the oven is shown in Figure 2.13 (a).

The temperature is an important parameter in the MIC process. For that reason, we calibrate the offset between the temperature measured by the internal thermocouples (oven setpoint) and the real temperature. To this end, an external thermocouple with tolerence of 2 °C is introduced and placed at the sample position in the tube. When the temperature is stabilized on both the oven display and the external thermocouple, we record the two temperatures in order to evaluate the offset. We measured the temperature by steps of 50 °C from 200 °C to 400 °C. Figure 2.13 (b) shows the measured temperature as a function of the oven setpoint and the corresponding offset (Figure 2.13 (c)). The oven always overestimates the process temperature by an amount that varies slightly with the temperature. For 300 °C, we estimate an offset of about 36 °C.

This calibration has to be done each time the internal thermocouples are moved (for instance during oven maintenance). The offset can vary by tens of degrees for a slight change in the position of the internal thermocouples.

Chemical etching

Once the annealing is complete, it is necessary to chemically remove the metal to expose the Ge platelets. Among many possible Al wet etch solutions, we find that only two can etch Al selectively with respect to Ge, namely H 3 PO 4 and HF. In this study, we usually etch samples in H 3 PO 4 heated at around 50-60 °C for a few tens of seconds until the metal is completely removed (as signalled by a change of color of the sample). Some samples may require slightly longer etching time depending on the temperature of the solution, which is not precisely controlled. Alternatively, the samples can be etched in diluted HF (1%-5%), typically during 20 s for an Al thickness of 20 nm. Increasing the HF etching time causes the etching of the silica under the Ge layer. The use of HF is subject to a strict security protocol and we prefer using orthophosphoric acid for safety reasons.

Experimental results

As seen before, many experimental parameters can affect the behaviour of the MIC process and the final characteristics of the crystalline Ge layer.

In the following sections we investigate the respective effects of:

• the thickness of the Al and Ge layers (in the range 10-40 nm)

• the ratio of the thicknesses of the Al and Ge layers

• the nature of the oxide interlayer

• the thickness of the Ge underlayer

Differential Interference Contrast microscopy (Nomarski)

In this section, we show how the samples can be characterized using Differential Interference Contrast (DIC) microscopy1 , which gives information on the topography of the surface by transforming optical path variations in the sample into amplitude variations.

Since we only work with opaque specimens, the microscope operates via reflection. With DIC, any optical path gradient in the specimen (induced by spatial variations of height or nature of material) is transformed into an image contrast. The light is generated by a halogen lamp and then polarized (Figure 2.14). The linearly polarized light is then reflected in the direction of the microscope frame by a half-mirror positioned at 45°. It then falls and meets the Nomarski prism that separates the beam into two coherent orthogonally polarized beams, named ordinary and extraordinary, with the same amplitude. The two beams are directed toward the condenser lens, that in a reflective DIC microscopy is also the objective. They are spatially displaced on the specimen by a fraction of micrometer. After been reflected, the light is collected by the objective and passes again through the Nomarski prism. The optical path difference of the sample is analyzed through the interference between the two orthogonal wavefronts.

Optical images formed with standard optics (bright field mode) and Nomarski optics can be seen in Figure 2.15 (a) and (b) respectively. For our samples, Nomarski images are equivalent to standard bright field images but with a tunable contrast, which can be adapted to highlight specific features.

The influence of the thickness of the layers

We investigate the influence of the thickness of the semiconductor and metal layers on the morphology of the final crystallized Ge layer, in terms of surface coverage, roughness, grain size and density. We use Nomarski optical microscopy to evaluate these parameters. We start by jointly varying the Al and Ge layers thickness, keeping a constant Al:Ge thickness ratio, and then we keep one thickness fixed and vary the other.

Constant Al:Ge thickness ratio

We studied three samples with different Al and Ge thicknesses but with about the same 1:1 Al:Ge thickness ratio. Details of the samples are given in Table 2.1. The samples have a Ge interlayer of 2 nm, oxidized in air for 9 min and no Ge underlayer. They are annealed in the trizone oven at 255 °C for 60 h.

Nomarski images of the samples after crystallization and Al wet etching of the top Al layer are shown in Figure 2.16. Small grains were formed in sample #2, with intermediate thickness. The thicker sample #3 exhibits larger grains, but it also shows a rough surface (dark), of unknown nature, between the Ge crystals. The thinner sample #1 shows the best morphology with large flat Ge islands (∼20 microns wide).

#1 (13E3) #2 (21L1) #3 (21G1)

Conclusion

The experiment is only partially conclusive. Different thicknesses of the layers could require separate optimization of the annealing conditions. Since we do not observe a real advantage in increasing the layer thickness, we preferably focus on the thinner ∼20 nm layers, since it requires less material.

In the absence of a clear effect, we preferably use Al/Ge structures with thin layers (∼ 20 nm). We now keep the Ge layer thickness constant and we vary the Al layer thickness. We estimated the Ge thickness to be about 20 nm. We fabricate three samples, detailed in Table 2.2, with three different deposited Al thickness of ∼15 nm, ∼20 nm and ∼40 nm.

Standard optical characterization after annealing is shown in Figure 2.17. The sample are not etched, so that we see the top layer, which should be Al if the layers correctly switched position during the MIC process. We observe that:

• With less Al than Ge (sample #1 in Figure 2.17), a large fraction of the top surface is dark. This reflects a large presence of Ge, since Ge is less reflective than Al.

• When the thickness of Ge and Al are comparable (sample #2 in Figure 2.17) the density of the dark areas, as well as their size, has significantly decreased.

• With more Al than Ge, we observe more complex features. Possibly, a Ge-rich material replaces the Al layer at the top of the sample. We did not further investigate this sample, due to the poor surface condition.

With Ge layer fixed at ∼20 nm, there is no improvement gained for a thinner or thicker Al layer. We now keep the Al layer thickness constant and we vary slightly the Ge layer thickness, around the optimum condition found previously. For an Al layer thickness of ∼20 nm, we fabricate three samples with Ge thicknesses of ∼16 nm, ∼20 nm and ∼24 nm. Details of the samples are given in Table 2.3. Nomarski images of the samples after etching the Al top layer are shown in Figure 2.18. We observe that:

• For Ge layer thinner than Al, the coverage of the silica is not complete so that it can be seen through the holes in the Ge layer (darker areas in Figure 2.18 (d)). • With a ratio Al:Ge ∼1:1, we partially fill the gaps in the Ge layer (Figure 2.18 (e)).

• At larger Ge thickness, the silica is completely covered, but the Ge layer is rougher (light grey areas in Figure 2.18 (f)).

An Al:Ge thickness ratio close to ∼1:1 with individual thickness of about ∼20 nm gives the best results. We choose this configuration for future optimizations.

The role of the oxide interlayer

Description of the experiment The interlayer plays a crucial role in the MIC process. As seen in section 2.1.3.2 p. 12, this thin layer is made of a material in which Ge has a low diffusion rate. It acts as a barrier to control the diffusion of Ge inside Al. We use AlO x or AlGeO x formed by exposing to air the Al layer or a thin Ge layer.

Toko et al. [Toko, 2012;Toko, 2014a] report that the crystallization rate, as well as the final Ge crystalline texture depends on the type of oxide and the oxidation time of the interlayer. Therefore, we investigate the effect of different interlayer oxides in order to find the best configuration. We compare samples with no oxide, with an AlO x oxide and with different AlGeO x oxides.

#1 (50C2) #2 (50R4) #3 (64I4)

Results

Sample #1 in Table 2.4 and Figure 2.19 has an AlO x interlayer, formed by oxidizing the Al (for 21 min) before depositing the a-Ge on top. Figure 2.19 (a) shows the sample after 20 h of annealing at temperature up to 350 °C; the Al has not been etched. We observe some dendrites, which could be crystallized Ge, and some brighter areas which could be the inverted Al layer. However, most of the sample appears as when initially deposited (i.e no layer inversion is observed).

For samples #2 to #4 in Table 2.4 and Figure 2.19, the interlayer is formed by oxidizing the Al layer protected by a thin amorphous Ge layer of 1.1, 1.5 and 3 nm thickness, respectively. The oxidation is performed in air for 12 to 43 min.

Nomarski optical characterization shows the samples after annealing for 20-30 h and after chemical wet etch of the Al layers. Sample #4 (Figure 2.19 (d)) shows smaller grains, compared to samples #2 and #3, with respectively 1.1 and 1.5 nm Ge interlayers, which present the best morphology. For the details of these samples, see Table 2.4.

#1 (50C2) #3 (64I4) #5 (32P1) #2 (50R4) #4 (50A2) 20 μm (a) (b) (c) (d) (e)
Sample #5 is fabricated without breaking the vacuum between Al deposition and a-Ge deposition (Figure 2.19 (e)). In this case, there is no interlayer and consequently we performed a very fast annealing (less than 1 h at < 250 °C). The layers do not exchange positions as in normal MIC. Instead, Al and Ge grains alternate in a single layer. This creates very small Ge grains of double thickness compared to that of the initial Ge layer.

Discussion If the barrier is too large, the Ge atoms do not diffuse through it and the MIC process is inhibited. This is clearly the case of sample #1, for which only a very limited fraction of the surface evolves during annealing, despite the high temperature (up to 350°C) used. Our results with a pure AlO x interlayer does not match the observations of Toko et al. [Toko, 2012;Toko, 2014a] who report fully inverted layers, but using thicker Ge and Al layers (50 nm).

Without any diffusion barrier (sample #5) the crystallization is very quick (< 1 h). However, it does not produce the expected layer exchange and produces very small grains. This shows that there are other possible crystallization mechanisms and that the interlayer needs to be tailored to favor the result of interest.

In samples #2 to #4, with an interlayer made of mixed Al-Ge oxides (AlGeO x ), the a-Ge crystallize following the classical layer exchange mechanism. This indicates that our formed AlGeO x interlayer is more permeable to Ge atoms than our pure AlO x interlayer (sample #1).

As we increase the Ge coverage of the Al layer, we observe that we can reduce the annealing temperature and duration to obtain a complete inversion of the layer. Since the GeO x oxide forms more slowly (a few hours) [Sahari, 2011] than AlO x (a few seconds), we suggest that the deposited Ge layer is not continuous but that it gradually covers the surface of the exposed Al layer with increasing thickness. Therefore a thicker Ge interlayer does not produce an AlGeO x interlayer richer in GeO x , but rather provides an increased coverage of the Al layer, which in turn reduces the surface of AlO x formed when the later is exposed to air. For 3 nm Ge deposition (sample #4) we propose that Ge fully covers the Al layer and that the final interlayer may only consist of GeO x .

These conclusions are summarized in Figure 2.20.

Conclusion

The deposition of a thin Ge layer over the Al before exposure to air permits the crystallization of the a-Ge following the standard layer exchange mechanism. In this way, we are able to crystallize thin Ge layers (20 nm), in conditions (24 h annealing at 250-300 °C.) similar to those reported in the literature for thicker layers (50 nm) using a pure (Ge-free) AlO x interlayer. In the following, this air-oxidized 1.1 or 1.5 nm Ge interlayer will be used in most of the samples.

• In contrast to the existing literature, we find that the direct oxidation of the Al layer in air creates an unsuitably strong AlO x interlayer, which inhibits the layer exchange mechanism (in our experimental conditions).

• The air-oxidation of the Al layer, partially covered by a thin Ge layer (1.1-1.5 nm) allows us to recover the standard MIC process. Slightly higher annealing temperatures or longer durations are required for the thinner (1.1 nm) compared to the thicker (1.5 nm) Ge interlayer.

Impact of the germanium underlayer

Description of the experiment Most of the samples studied in the previous section had a thin Ge underlayer (1 nm) between the silica substrate and the Al layer. This Ge layer is expected to act as a "wetting layer" [Chen, 2010] i.e. change the surface energy, for the subsequent Al layer deposition. It also has a large impact on the MIC crystallization temperature [Numata, 2014a]. We investigate here how the morphology of the final Ge platelets change as a function of the thickness of the Ge underlayer.

#

Results

The characteristics of three samples, respectively without the Ge underlayer, and with 1 nm and 2 nm Ge underlayers, are reported on Table 2.5. Nomarski optical images taken after etching of the Al layer are shown in Figure 2.21.

Sample #2 (Figure 2.21 (b)) appears to have the largest and flattest grains. When we increase the Ge underlayer thickness, sample #3 appears to have small grains and a stronger double layer roughness.

Although we annealed all samples together, sample #1 (without underlayer) seems to have badly crystallized areas.

Discussion

The use of the Ge underlayer is critical to obtain wide flat Ge platelets, at least in our experimental conditions. The difference in morphology without (sample #1) and with (sample #2) is enormous and highlights the difficulties in optimizing the sample structure. The Ge underlayer is positioned between the Al layer and the silica surface, right at where the migrating Ge atoms are supposed to nucleate in the layer exchange mechanism. This pre-saturation of Al-silica interface is very sensitive to the amount of Ge and the over-saturation of the interface (sample #3) leads to a completely different morphology. Note that the optimization of the Ge underlayer can only be made in combination with that of the (Ge)AlO x interlayer (see previous section 2. 3.3).

The crystallization process at play without underlayer (sample #1) looks similar to the expected Al induced crystallization of the Ge layer but it is not clear how the Ge dendrites relate to the large Ge platelets obtained with a Ge underlayer. The optical in-situ characterization performed later (see section 2.4.2.2) will provide more details on dendrite formation.

The Ge underlayer is critical to obtain large and flat Ge platelet by the layer exchange mechanism.

Conclusion

In this section, we have determined the experimental parameters for the Metal Induced Crystallization of amorphous Ge in the presence of Al focusing on thin layers.

Compared to the oversimplified initial scheme of the layer exchange mechanism, in which the amorphous Ge simply exchanges position with the Al layer as it crystallizes, we find that the sample structure required to activate this process is much more complex than a simple metal-semiconductor bilayer.

The introduction of an interlayer between the Al and Ge layers is necessary to regulate the Ge diffusion so that the Ge crystal nucleation takes place at the interface between the Al and the silica substrate. In addition to the interlayer, we need to add a Ge "underlayer" between the Al and the silica, which pre-saturates the Al-silica interface in Ge, to finally obtain the expected large Ge grains (50 µm wide) instead of irregular Ge dendrites.

Although the literature describes issues when crystallizing Ge layers less than 50 nm thick, we show here that thin Ge layers (∼20 nm) can be crystallized if the interlayer and underlayer are carefully tuned: we actually need to introduce a mixed (Ge,Al) oxide interlayer instead of the pure AlO x and to use a thickness of 1 nm for the Ge underlayer.

Despite the lower layer thickness (20 nm) used in our experiment, we find the same optimum of 1:1 thickness ratio between the main Al and Ge layers than reported in the literature for thicker structures (50 nm thick layers).

For a given sample, the annealing temperature and duration play an important role. However the optimum annealing conditions heavily depend on the structure of the sample. Without an interlayer, Ge can crystallize in less than an hour at 250°C, whereas with a Ge-free AlO x interlayer, we never observe complete Ge crystallization, even using annealing temperature above 350°C.

Because each structure requires specific annealing time optimization, experiments based on fixed annealing duration and subsequent ex situ observations are tedious and time consuming. That is why we developed in situ observation during the annealing, as described in the next section. This is critical to measure the crystallization kinetics and identify the precise chronology of the MIC process.

In situ characterization and kinetic study

Monitoring and quantifying the density of crystalline nuclei and the crystallization rate in situ during annealing is highly preferable to performing time series in which the thermal treatments are stopped and resumed at regular intervals. For the optimization of (111)-oriented AIC-Ge, for which annealing durations can reach several tens of hours, in situ monitoring becomes a critical requirement. However, the technical difficulties in observing hot samples (200-350 °C) in vacuum or nitrogen ambient at high magnification have limited the use of in situ optical microscopy. Some recent works on AIC-Si are still based on low-resolution imaging [Kurosawa, 2014;Antesberger, 2012;Jung, 2010]. It is only in 2016 that Tutashkonko et al. [START_REF] Tutashkonko | Effects of the Si/Al layer thickness on the continuity, crystalline orientation and the growth kinetics of the poly-Si thin films formed by aluminum-induced crystallization[END_REF] reported better resolution (5 -10µm) which allowed these authors to better characterize the processes at play during the AIC of Si. Here, a specially designed in situ optical monitoring setups allows us to characterize the AIC of amorphous Ge with an even higher spatial resolution (∼ 1 µm), close to the optical limit.

In the following section, we first describe the in situ imaging setup, then we establish a chronology of the different MIC processes and we quantify their respective kinetics.

Assembling the in situ microscope

Oven and optics

We use a long working distance microscope 50x objective (Mitutoyo Plan Apo 378-805-3), with an optical resolution limit of 500 nm, associated with a programmable hot plate (HARRY GESTIGKEIT) which lid has been extensively modified to provide a uniform nitrogen supply and accommodate an optical viewport made of fused silica (Figure 2.22). The optical setup is composed of a CCD camera, an optical tube, a connector for the lamp and for the 50 X objectives (Mitutoyo). The optical tube is clamped to a micrometric adjustment plate which permits up and down movements of the whole optics to focus the image. The movement is actuated by a motor controlled by USB. The whole optical setup is mechanically fixed to an optical table. The light is injected sideways in the optical tube and it is focused by the objective on the sample (in-lens illumination). A flux of nitrogen is flown in the closed chamber (at atmospheric pressure) to limit the oxidation of the surface during the annealing. The flow rate of nitrogen is controlled with a manual valve.

Thermocouple

Precise temperature measurements of the sample surface are obtained using an external thermocouple introduced through the lid of the oven. An example of temperature monitoring is shown in Figure 2.23. Temperature recording and image acquisition are controlled via the same software (Sharpcap, version 2.9.3030) using a specially written Python script.

Autofocus

After damping the mechanical vibrations of the whole system, the optical resolution is ultimately limited by the turbulent air flow between the relatively hot fused silica viewport and the cooler surface of the microscope objective. At such a high magnification, we find that a motorized automatic focus is necessary to maintain sharp images, given the different thermal expansions of the various elements during tens of hours of annealing. An example of optical monitoring without the autofocus is shown in Figure 2.24 (a-c) along with the corresponding temperature measurement (d). The sample is manually kept in focus during the first 45 min until the temperature stabilizes at 300 °C. After 15 min, the sample starts to be out of focus (Figure 2.24 (b)) and it is completely out of focus after 45 min (Figure 2.24 (c)).

For this reason, we set up an autofocus system which checks the image and actuates the motor to obtain the best focus. The motor is mechanically coupled to the finger- In details, the software computes the contrast of a selected region. The motor is actuated in forward and reverse directions and the changes in contrast are monitored by the software. When the best focus position is found, the image is saved. At the same time, the temperature reading from the thermocouple and the elapsed time are also saved. After a fixed waiting time, a new image is acquired. This sequence is summarized in Figure 2.25.

Investigating MIC processes using in situ microscopy

With the setup described above, it is possible to follow the evolution of the MIC process in real time, adjust the temperature and stop the annealing as soon as the crystallization is completed.

Typical MIC experiment -bottom view

Description of the experiment In this experiment, we use in situ optical microscopy to follow the MIC of a typical sample, in order to gain a general overview of the process. We first use a thermal silica substrate in order to characterize the reaction at the SiO 2 /Al interface, i.e. viewing through the transparent silica substrate. We place an additional piece of silica on top of the metallic heater, to prevent contamination of the sample and to create conditions similar to the standard MIC, where the heat flows through the silica substrate (Figure 2.26). The light goes through the transparent substrate, illuminates the interface between Al and silica and is finally reflected to the microscope.

Metallic support

Results

Figure 2.27 illustrates the optical monitoring of a sample during annealing. After an initial temperature ramp from ambient to 238 °C (∼1 h), the sample is maintained at a fixed temperature (238 ± 1 °C) for 16 hours while optical images are acquired every 30 s. We monitor the optical changes in the Al layer at the bottom interface. After 3 min (Figure 2.27 (a)), we observe the formation of small spots which then grow in a dendritic fashion over several tens of microns (arrows, Figure 2.27 (b)). We refer to this process as MIC-1 in the following. After about 2.9 h of annealing, a new darker contrast appears near the previously formed dendrites (red arrows, Figure 2.27 (f) and Figure 2.27 (q)). This feature then slowly expands, encircling the previous dendrites (Figure 2.27 (h)) and merging with other darker fronts (Figure 2.27 (l-m)) until it finally fills the field of view in approximately 14 h (Figure 2.27 (n)). We will call this second process MIC-2.

High magnification images (Figure 2.27 (p) and Figure 2.27 (q)) show that both processes are actually dendritic. MIC-1 (Figure 2.27 (q), gray contrast) produces long dendritic branches which are typically 20-30 microns long and 2 µm wide. MIC-2 (Figure 2.27 (p), darker contrast) yelds a dense collection of small dendrites, with a typical scale of 1-2 microns.

At this stage, the simple observation of dendrites is not enough to ascertain the crystallinity of the final Ge platelets. However, we will prove in section 2.6 p. 54, that these structures are actually Ge crystals and that the two processes, MIC-1 and MIC-2, are indeed related to the Al-induced crystallization of the amorphous Ge. To simplify the discussion, we call the progress of the MIC-1 and MIC-2 processes 'crystallization rate' and 'crystallization kinetics' in the following sections.

Optical in situ microscopy allows us to distinguish within AIC two different processes that operate sequentially, which we call MIC-1 and MIC-2.

The observation of two distinct phenomena raises the question of how they relate to the single layer exchange mechanism described in the literature.

Understanding the effect of the germanium underlayer

Description of the experiment Samples with and without the a-Ge underlayer result in very different morphologies for the final Ge platelets (see section 2.3.4 p. 33) We now investigate how the absence of underlayer affects the kinetics of the previously identified MIC1 and MIC-2 processes. After an initial temperature ramp from ambient to 255 °C (∼1 h), the sample is maintained at this temperature (within ± 2 °C). MIC-1 dendrites appear after 0.6 h (arrows, Figure 2.28 (a)) similarly to what was observed in Figure 2.27 (ab) with the Ge underlayer. The branches of the dendrites first grow in length (Figure 2.28 (b) and Figure 2.28 (e)) and then stop (Figures 2.28 (c) and Figure 2.28 (f)). Then, a second growth process occurs (arrow, Figure 2.28 (d) and Figure 2.28 (g)) after 2.45 h annealing, producing darker zones that extends, starting from the previously formed MIC-1 dendrites. This is the previously identified MIC-2 process. A schematics differentiating MIC-1 and MIC-2, relatively to Figure 2.28 (g), is shown in Figure 2.28 (h).

Results

Discussion

In this experiment, the Al layer is deposited directly on the silica substrate (see insert in Figure 2.28 (a)). Apart from this, the rest of the sample structure is the same as described in Figure 2.27, so that the only difference is the absence of the Ge underlayer.

We observe the two same crystallization phenomena, namely MIC-1 and MIC-2, with and without underlayer. However, the sample without underlayer, (Figure 2.28), shows a larger extension of the MIC-1 dendrites whereas MIC-2 (darker area) fails to cover the rest of the surface and simply stops progressing at 250°C. The opposite configuration, large MIC-2 area, is found in the sample with a Ge interlayer.

Without underlayer, a further increase of the temperature to ∼325 °C only slightly increases the final MIC-2 coverage (2.28 (i)) but large areas still remain unaffected. This crystallization problem correlates with the poor morphology previously observed in section 2.3.4 p. 33, for the sample with no Ge underlayer.

Using the in-situ monitoring, we indentify the two crystallization processes MIC-1 and MIC-2, independently of the presence of the Ge underlayer. The difference in the final morphology of the Ge platelets without Ge underlayer is linked to the spatial extension of the MIC-1 relative to the MIC-2 process.

Crystallization kinetics of the MIC-1 and MIC-2 processes

Up to now, the in situ observation allowed us to visually recognize two process during the annealing, which we called MIC-1 and MIC-2. We now characterize the kinetics of these two processes using the previously acquired in situ movies.

Bottom view: through the substrate

To better appreciate the different kinetics of the two phenomena occurring at the same temperature (238 ± 1 °C), we quantify the spatial extension rate of MIC-1 and MIC-2 using the frames from Figure 2.27 (section 2.4.2.1).

In order to do so, we automatize the image processing using software ImageJ (or Fiji)2 . Figure 2.29 (a) shows an image before any processing. All the images are first converted to grayscale and then aligned with each other3 , Figure 2.29 (b). A simple value threshold is then set to measure the extent of the features on a predefined profile line (Figure 2.29 (c)). The profile lines are chosen along the center of the dendrite branches, from the dendrites shapes in the final image.

Figure 2.30 shows the spatial extension of large dendrites created by the MIC-1 process. We find that the 'arms' of the MIC-1 dendrites all expand at a fixed rate of 10.5 ± 0.3 nm/s until they stop completely. The stopping time can vary between dendrites and between the arms of the same dendrite. All the observed MIC-1 dendrites stop growing after 1 h or 1.5 h of annealing, well before reaching another MIC-1 dendrite.

The situation is very different for the MIC-2 process (Figure 2.31). It develops at a slower but steady growth rate of 2.1 nm/s, measured over more than 7 h. The crystallization front progresses in an isotropic manner, with the same velocity for all dendrites, despite possible different starting times. In the presence of a Ge underlayer, the MIC-2 process only stops when the whole surface of the sample has been covered or if the sample is cooled down before that. The crystallization rate of MIC-1 is 5 times faster than that of MIC-2, for the same temperature (Figure 2.31). The MIC-1 process stops, for some unknown reason, well before the complete coverage of the surface.

Top view (a-Ge/air interface)

Metallic support light a-Ge Al SiO 2 /Si substrate

Description of experiment

We now characterize the sample from the top (a-Ge/air interface) during the annealing. This allows us to link the previous in situ observation from the back (section 2.4.2, Figures 2.27) to standard microscopy observation from the top performed after the annealing.

The sample is positioned directly on the heating metallic support (Figure 2.32). The layer structure is the same as for the samples in section 2.4.2.1, but we use thermal silica grown on a silicon substrate since we do not need a transparent substrate. We repeat the kinetic measurements of section 2.4. 3.1 (illustrated in Figures 2.30 and 2.31) in this new configuration. Results Figure 2.33 (a) gives the variation of the length of three MIC-1 dendrite branches, measured from the center of the dendrites, as a function of the annealing time at ∼237 (constant within ± 2 °C). As already observed, the MIC-1 dendrites first grow at a uniform rate but they stop at different times for each dendrite. This first expansion rate equals 6.8 ± 0.3 nm/s, as estimated from a linear fit of the first part of the graph. 2.33 (b) shows how the expansion length varies with time in the MIC-2 process. This process does not consist in individual dendrites but rather produces uniform light areas. Since these areas are somewhat irregular, we estimate the MIC-2 extension either by direct linear measurements (L1 and L2 in the figure) or as the radius of the disk having the same area as the MIC-2 zone. Both measurements give a rate of 2.5 nm/s over 10 h annealing at 237± 2 °C, is. As previously observed from the bottom, the MIC-2 areas uniformly expand and merge to cover the entire surface of the sample.

From in situ top views, we identify the same MIC-1 and the MIC-2 as viewed from the bottom of the structure. The MIC-2 appears as a compact growing disk instead of an ensemble of dendrites. The kinetics measured from the top confirm that measured from the bottom.

Conclusion

From the ex-situ optimization of the MIC structure (section 2.3), it appears that the outcome of the MIC is so varied that it could involve more than one crystallization process.

The optical in situ characterization during annealing confirms this hypothesis and allows us to establish a precise chronology of the crystallization sequence and we identify at least two distinct crystallization processes, MIC-1 and MIC-2.

Using transparent substrates, high resolution optical images show that both MIC-1 MIC-2 proceed by dendritic growth. MIC-1 is responsible for the formation of large Ge dendrites, visible from the front and the back fo the sample, which disrupt the morphology of the final layer in the absence of a Ge underlayer. MIC-2 only forms small dendrites, visible only from the back of the sample, which are responsible for the formation of large Ge platelets.

The two phenomena have different timings and crystallization kinetics:

• MIC-1 starts first and forms large dendrites which branches expand at constant linear rate, before they slowly stop growing;

• MIC-2 starts after MIC-1 and generates a disk which radially expands until all the sample has reacted.

In the next section, we address the crystallinity and the crystal orientation of the material created by these two MIC processes.

Chemical analysis of MIC-1 and MIC-2 products

The MIC-2 is the process responsible for the formation of the large flat Ge platelets.

The optical images suggest that MIC-2 proceeds by layer inversion, due to the strong change of contrast.

In this section, we investigate the features created by MIC-1 and MIC-2, before and after the annealing, to determine the respective growth mechanisms.

Qualitative analysis by SEM-EDXS

Description of the experiment

To better characterize the MIC processes, we partially anneal (6 h at 260 °C) a typical sample (Ge/Al/GeO x /Ge 1/20/1.5/20 nm) and we characterize it from the top by Nomarski imaging, scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy performed in the SEM (SEM-EDXS).

To identify the location of Al and Ge species, we perform EDXS experiments at two different acceleration voltages, 2 kV and 7 kV. At 7 kV, the electrons traverse the whole Al/Ge bilayer, while at 2 kV the signal is more sensitive to the top layer (see Monte Carlo simulation in next section). This enables us to probe the sample structure at different depths.

Results Figure 2.34 (a) shows Nomarski images of the partially annealed sample in a region where both the MIC-1 dendrite, the MIC-2 area and the reference structure co-exist.

The SEM image and the corresponding and EDXS map of a similar area are shown in Figure 2.34 (b) and (c) for an acceleration voltage of 2 kV and in Figure 2.34 (d) and (e) for an acceleration voltage of 7 kV.

In Figure 2.34 (d), we identify four different zones, marked by squares E0, E1, E2 and E3

• E0 is the reference structure, not yet affected by either MIC-1 or MIC-2.

Consequently, at 7 kV, both Al and Ge are detected (Figure 2.34 (f)) whereas the Ge signal of the amorphous top layer predominates at 2 kV (Figure2.34 (c)).

• E1 is the dendritic area created by MIC-1; the Ge signal dominates at both accelerating voltages.

• E2 is the large area created by the MIC-2 process. The layer consists of Al and Ge (as shown by the 7 kV image) but the surface is Al-rich, which suggest that layer inversion has occurred.

• E3 is a zone at the edge of the area affected by MIC-2. The Al signal dominates at the surface and the full structure.

Our qualitative EDXS analysis strongly suggests that layer exchange occurs at the MIC-2 stage, when the surface changes from Ge-rich to Al-rich between the MIC-1 and MIC-2 zones.

The structure of the MIC-1 area remains unclear, except that it contains more Ge that the rest of the layer. An Al-rich zone is observed at the edge of the MIC-2 area.

Quantitative analysis by SEM-EDXS

Descritption of the experiment To better evaluate the different amounts of Ge and Al, we extract the EDXS signal from areas identified in section 2.5.1 (Figure

(d))

. We obtain a value for the amount of Ge and the Al by integrating the area below the respective EDXS peaks after we subtract the background by polynomial fitting (Figure 2.35).

We use the software CASINO (monte CArlo SImulation of electroN trajectory in sOlids) to simulate the EDXS signal from Al and Ge layers of various thicknesses at different primary electron energies. Discussion While it is not always possible to find a perfect match over the whole electron energy range considered, we can reproduce rather satisfactorily the measurements carried out in each zone to a likely structure consisting of two or three layers of different thicknesses. These best-fitting couples are as follows.

Results

2.35:

• As expected, E0 corresponds to the as-deposited structure: equal thicknesses of Al and Ge, with the latter on top.

• E1, created by MIC-1, consists of a thick Ge layer with a thin underlying Al layer.

• E2, created by MIC-2, may fit with a three-layer structure: Al on the top, then Ge and possibly Al at the bottom.

• E3 is the most problematic zone and loosely fit with a pure Al with traces of Ge.

The quantitative analysis only brings marginal improvement over the previous qualitative observation. The only clear match is the reference area (E0)

, not yet affected by the MIC-1 and MIC-2. The other areas (E1-3) fit with different layer structures and present too much signal intermixing to unambiguously determine the layer structure.

Analysis of the MIC-1 and MIC-2 areas by crosssectional TEM

Cross sectional transmission electron microscopy (TEM) is an ideal technique to fully characterize how the MIC-1 and MIC-2 processes affect the initial layer.

Description of the experiment

To make sure that TEM analysis is performed on a relevant area, we use a focused ion beam (FIB) to extract a TEM foil from a zone previously analyzed by top view SEM-EDXS. In this way, we can link the large scale EDXS analysis of the previous section to the present detailed TEM observations.

Results

We first identify the reference area (unreacted, R label, dark gray) using the SEM secondary electron (SE) image (Figure 2.36 (a)). The contrast in this SE image is not high enough to distinguish between the MIC-1 and the MIC-2 zones (both light gray) but the corresponding EDXS image ( To help visualizing the MIC-1 area in the SE image (Figure 2.36 (a)), we draw a thin white line delineating the feature observed in the EDXS map (Figure 2.36 (b)). We then investigate the interfaces between the different zones by cross-sectional TEM. The sample is a slice extracted by FIB along the dashed line of Figures 2.36 Two areas are of particular interest, the R/M2 interface at the crystallization front between MIC-2 and the reference area (Figure 2.36 (c) to (e) ), and region M2/M1 at the interface between MIC-1 and MIC-2 (Figure 2.36 (f) to (h)). The R/M2 region shows the inversion of the Ge layer, initially at the top (in the R reference area) and then at the bottom of the stack (zone M2), after the MIC-2 process, with a mixed (Al and Ge) region at the growth front (i). The region M2/M1 shows a different structure with the inverted layering in the M2 area (Ge below Al) and a double layer of Ge in the MIC-1 area, without any Al (Figure 2.36 (f,g)).

20 μm R/M2 (a) (b) M2/M1 R/M2 M2/M1 (f) (g) (h) (c) (d) (e) R M2 M2 M1 i i I R/M2 M2/M1 R M2 M1
Both areas R/M2 and M2/M1 yield similar oxygen signals. In each of the two areas we distinguish the silica substrate, the AlGeO x interlayer and top surface oxidation (white arrows, Figure 2.36 (e) andFigure 2.36 (h)).

Discussion

It is now clear that the MIC-2 process is indeed the standard alumininium induced layer exchange (ALILE) crystallization mechanism, as initially reported on amorphous Si [Nast, 2000b;Wang, 2007] and later on amorphous Ge [Kurosawa, 2012;Nakazawa, 2013b]. In the TEM images, the Al and Ge layers are clearly seen to exchange positions. This observation confirms the previous observation by top view SEM-EDX (sections 2.5.1 and 2.5.2)

Let us now discuss the M2/M1 interface of our partially annealed sample (Figure 2.36 (a), blue rectangle, and Figure 2.36 (f) to (h)). Here, the standard structure produced by the ALILE process (MIC-2) terminates on a double layer of Ge, present in the M1 area. The top-view chemical image indicates that this double layer exists not only in the M2/M1 region, but also extends to the reference area yet unaffected by the ALILE (MIC-2) process (Figure 2.36 (b), bottom part of the image).

Combining all ex-situ measurements, TEM-EDX and SEM-EDX, we can conclude that the large MIC-1 dendrites consist of a double layer of Ge. The later is created prior to the ALILE (MIC-2) process, since MIC-2 has not yet reached this part of the sample. This observation is consistent with in situ optical microscopy, which discriminates two phenomena, the ALILE (MIC-2) and the MIC-1, with different timings and kinetics (see section 2.4). Putting this together, we can conclude that the large dendrite with a double layer of Ge are created by the MIC-1 process.

A careful reading of the literature suggests that MIC-1 might already have been observed, but not identified as a distinct process. Toko et al [Kurosawa, 2009b;Toko, 2013a] observe that double layers of Ge can form during the Ge-AIC.

AlGeO x a-Ge SiO 2 Al Ge Al ALILE / MIC-2 MIC-1 a-Ge SiO 2 Al Ge Al Ge ? (a) (b)
a-Ge? c-Ge? c-Al? c-Ge? Both processes start from the same structure, the Al/Ge bilayer, operate via the same mechanism of diffusion of Ge into the Al layer but assume different kinetics (see section 2.4.3) and produce different final structures.

By combining in-situ monitoring, SEM EDXS and cross-sectional TEM, we unambiguously determine the sample structure after the MIC-1 and MIC-2 processes:

• MIC-1 is a novel crystallization mechanism, which creates a double Ge layer;

• MIC-2 is the standard ALILE process, producing the usual exchange of positions between the Al and Ge layers..

Conclusion

In this section we have used a wide range of tools and techniques to establish the nature and the characteristics of the processes at play during the annealing of our samples. Our original in situ monitoring revealed not one but two distinct crystallization mechanisms, with different timings and kinetics. Top-view SEM-EDXS shows that the surfaces resulting from these two processes are chemically different but cross sectional TEM was required to unambiguously determine the full layer structure.

We are able to confirm that MIC-2, which creates the large platelets, is the standard Aluminium Induced Layer Exchange (ALILE) process. Conversely, MIC-1 was not described before, or possibly fortuitously ascribed to a roughening of the surface. We show here that it is a parasitic process, which occurs before the ALILE, at lower temperature, and which creates a bilayer of pure Ge instead of the expected Al/Ge structure.

Reference experiments carried out at 270 °C (not shown) indicate that, in the absence of Al, the Ge layer does not evolve within 40 h. Therefore, MIC-1 is also a genuine crystallization process, catalyzed by the Al layer.

Structural analysis of MIC-1 and MIC-2/ALILE

Grazing Incident X-ray Diffraction (GIXRD) and Electron Backscatter Diffraction (EBSD) can be used to assess the crystallinity of a sample and to measure the crystal orientation. In the following section, we investigate the crystal structure of the layers produced by the MIC-1 and the ALILE/MIC-2.

Analysis by in-plane X-ray Diffraction

We use GIXRD to evaluate the crystallinity and the crystallographic texture of samples after the MIC-1 and MIC-2/ALILE processes. The specific X-ray Diffraction (XRD) configuration is detailed in appendix A. The analysis of the MIC-1 dendrites requires a specific sample where we use the in situ optical monitoring to stop the annealing before the onset of the MIC-2/ALILE process. This allows one to selectively investigate the crystal orientation created by the MIC-1 mechanism.

In Figure 2.39, we compare GIXRD diffractograms of the unannealed reference sample A, of sample B with only MIC-1 dendrites and of fully annealed sample C with MIC-1 and MIC-2/ALILE structure covering the whole surface. The corresponding Nomarski optical micrographs are presented in Figure 2.38. A 0 is the reference diffractogram of a single Al layer (20 nm), as deposited.

Given the large difference of GIXRD signal between the samples, the count rates are plotted in log scale in Figure 2.39. The baseline of each diffractogram is indicated by a solid black line. The as-deposited unannealed sample A shows no proper diffraction peak. This is compatible with amorphous Ge and very small randomly oriented Al crystals, as confirmed by the GIXRD of the single Al layer (A 0 in Figure 2.38). On the contrary, sample B exhibits weak diffraction peaks whereas sample C produces the most intense peaks. Comparison to the reference Ge powder diffractogram (R) (ICDD card 00-004-0545) shows that the 220, 422 and 400 diffracted beams are more intense in our sample, which indicates a preferential crystal orientation of the crystallites. .

2θχ (°)

Log Intensity (a. u.)

1 1 1 3 1 1 3 3 1 5 1 1 4 4 0 4 0 0 2 2 0 4 2 2 (A) (B) (C) (R) Al Al (A 0 ) Al only Al+Ge MIC-1 MIC-2
Powder 111) texture. 220,311,222) GIXRD for the ( 111) and ( 100) crystal texture of the samples (B) and (C) studied in Figure 2.39 The texture analysis is performed using the Lotgering factor adapted to GIXRD (see appendix A). In Table 2.6, we compare the values of the corresponding factors L XRD confirms that both MIC-1 and MIC-2/ALILE produce a crystallization of the amorphous Ge layer. However:

• The MIC-1 crystals are poorly crystallized and possibly (100)textured.

• The MIC-2/ALILE crystals are highly crystallized and (111)textured.

• the Al evolve from small random crystallites to larger crystal.

Analysis by SEM-EBSD

Electron BackScatter Diffraction is a powerful tool that gives information about the crystal structure and the crystal orientation at each point sampled by an electron beam. An example of diffraction pattern with Kikuchi bands of a Ge crystal is shown in Figure 2.40 (b). The bands are automatically indexed by the software in order to extract the crystal structure and its orientation. If the structure is known, it is possible to extract the crystalline orientation from a minimum of three Kikuchi bands. In our case, Ge has the diamond crystal structure (see Figure 2.40 (c) and XRD result in Figure 2.39), so that we directly measure the crystalline orientation. It is then plotted as a color-coded map, as shown in Figure 2.40 (c). The EBSD analysis presented here has been performed in collaboration with Institut de Chimie Moléculaire et des Matériaux d'Orsay (ICMMO) using a SEM ZEISS Supra 55 VP equipped with a Hikari/OIM TSL EDAX detector.

Crystal orientation of the

-EBSD analysis of the Ge layer after complete AIC and removal of Al by wet etching (sample C of Table 2.6). The normal of the Ge crystallites is shown projected along the vertical direction (i.e. perpendicular to the substrate) in (a), and in-plane(parallel to the substrate) in (b).

Results

The EBSD maps are shown in Figure 2.41. Here, the top Al layer was etched chemically (in H 3 PO 4 ) to expose the bottom Ge layer, since EBSD is a surface characterization technique. Figure 2.41 (a) presents a map of the crystal normal projected along the vertical direction.

Ge crystals form isolated islands (around 30-40 µm in diameter), mostly (111)textured, each organized around a central dendritic zone with a typical (100) texture. As an example, in Figure 2.41 (a), a (111)-oriented island is delimited by a dotted line and the corresponding central dendrite by a full line. The in-plane projection of the crystallite normal is shown in Figure 2.41 (b), where the central 100 dendrites are marked by black lines. We observe that each large (111)-textured island is actually composed of different grains, with random in-plane orientations, organized around a dendritic (100)-textured core.

Discussion EBSD shows that the MIC-2/ALILE process, which results in coverage of the whole sample surface, yields mostly (111)-oriented Ge crystallites (Figure 2.41). This preferential (111) crystal orientation is also clearly visible in the GIXRD diffractogram and quantified by the calculated Lotgering values (section 2.6.1, p. 54). This type of (111) preferential orientation of the Ge crystallites has already been reported for similar samples in which a silica substrate, a Ge underlayer [Numata, 2014a;Toko, 2014a] and an AlGeO x [Numata, 2014b;Toko, 2014a;Hu, 2010] interlayer are combined and annealed at low temperature.

The dendritic shape of the (100) oriented regions visible in the EBSD maps (Figure 2.41), reminds one of the geometry observed optically in Figure 2.27, with a MIC-1 dendrite in the middle of a (111) island created by the MIC-2/ALILE process. This suggests that these partially (100)-oriented regions revealed by EBSD are related to the large dendrites created by the MIC-1 process. This hypothesis is also supported by the Lotgering values calculated from GIXRD (section 2.6.1) which hints at a preferential but partial (100)-orientation of the MIC-1 crystallites. Still, the unambiguous determination the crystal orientation created by the MIC-1 process require another specific experiment.

EBSD confirms that the large Ge platelets created by the MIC-2/ALILE process are (111) oriented crystals. Each large platelet is actually composed of several independent subgrains with (111) out-of-plane orientation and random in-plane orientation.

Crystal orientation of the top Al layer obtained by MIC-2/ALILE

Description of the experiment We now characterize the sample before the Al wet etching. We use a partially annealed sample so that zones with the MIC-2 and reference area coexist.

Results Figure 2.42 shows the selected area with the zone affected by the MIC-2/ALILE (left) and the reference zone (right). In Figure 2.42 (a) we see the Image Quality (IQ) map, which is a measure of the quality of EBSD pattern, where black corresponds to an absence of EBSD pattern. The MIC-2 area appears light grey, which indicates that measurable EBSD patterns were acquired in this zone, while the reference area is black, indicating that no EBSD pattern could be measured there. Figure 2.42 (b) shows the corresponding crystal orientation map along the direction perpendicular to the substrate. The EDXS signal, acquired at the same time as the EBSD, is used to draw the corresponding Al and Ge concentration maps (Figure 2.42 (c) and (d)). The chemical information from the unreacted zone is filtered out using the IQ signal (Figure 2.42 (a)).

Discussion XRD shows an Al signal much higher after the annealing than before (section 2.6.1). This indicates a recrystallization of the Al during the annealing. EBSD confirms that the top Al layer is crystalline but with a random orientation (Figure 2.42 (b)). The Al grain size is about 1 µm, which is much less that the lateral size of each sub-grain in the large Ge platelets (Figure 2.41). .

Despite the apparent symmetry in the layer exchange mechanism between the Al and Ge layers, the change of grain size and the difference in crystal texture indicate that the nucleation and the crystal growth are very different for each material. This correlates with the accepted mechanism in which the Ge first diffuses, nucleates and form growing crystal grains in the vicinity of the Al boundary; then, Al migrates (and recrystallizes).

During the MIC-2 /ALILE process, the amorphous Ge crystallize and the Al re-crystallize as they exchange position. While the bottom Ge layer shows large grains and a well-defined (111) crystal texture, the top Al layer presents smaller grains (∼1 µm) and a random crystal orientation. Results In Figure 2.43 (a), we first investigate the crystallinity of the top Ge layer. For this particular SEM observation, we use a specific imaging mode which is sensitive to the chemical nature of the full bi-layer. We use an acceleration voltage of 20 kV and a strong negative polarization of the SE detector to select backscattered electrons. In this way, we get a clear image of the dendrite, even if the associated topography is relatively flat.

Crystal orientation of Ge bi-layer obtained by MIC-1 Description of the experiment

Selected EBSD patterns of the top Ge layer inside and outside the dendritic area show no feature, which indicates that the top Ge layer is amorphous everywhere.

Then, we selectively wet etch (H 2 O 2 ) the top amorphous Ge layer to expose the bottom layer. Figure 2.43 (b) shows the SEM image (taken at 20 kV with standard imaging conditions for secondary electrons) along with the EBSD patterns of the Al layer (outside the dendrite area) and Ge area (inside the dendrites). We now observe a clear SE contrast between the Ge dendrite and the surrounding Al material.

The Al layer produces weak EBSD diffraction patterns as expected from the previous GIXRD (see Figure 2.39). The dendritic Ge displays strong diffraction patterns. A full EBSD analysis of the crystalline Ge area is shown in Figure 2.43 (c,d), using the EBSD image quality to filter out the signal from the Al layer. The Ge dendrites consist in small grains (a few micron wide) with a mixed ( 100) and ( 111) texture in the vertical direction (Figure 2.43 (c)), and a random in-plane orientation (Figure 2.43 (d)).

Discussion Double layers have also been observed during the AIC of SiGe material but with two superposed crystalline layers [Zhang, 2012]. Here the top Ge layer is largely unaffected by the MIC-1 process, which only leads to the crystallization of the bottom Ge layer, without pushing up the Al grains.

It is interesting to compare these findings with those illustrated in Figure 2.41, after both MIC-2/ALILE and Al wet etching. There, we can observe directly the mixed ( 100) and ( 111) crystal orientation of the lower part of the MIC-1 dendrite, which indicates that the amorphous top layer has disappeared prior to the EBSD characterization, either due to the chemical treatment or by diffusion during the rest of the annealing.

In all our samples, the MIC-2/ALILE process always starts in the vicinity of a MIC-1 dendrite. This suggests that MIC-1 creates boundaries with a high free energy, possibly strained or defective. The observation of mixed ( 100) and ( 111) texture in the MIC-1 areas indicates that Ge crystal nucleation may occur not only at the lower Al/SiO 2 interface [Kurosawa, 2012;Park, 2013] but elsewhere, possibly inside Al grain boundaries [Wang, 2008a;Zhang, 2012]. This suggests that the MIC-1 process may be related to small defect in the AlGeO x barrier that we are not able to resolve in by TEM (Figure 2.36). This hypothesis could also explain the different kinetics of the MIC-1 and MIC-2/ALILE processes: the first might involve nucleation inside the Al grain boundaries and the second nucleation at the Al/SiO 2 interface.

Combining EBSD with previous data, we obtain a clear description of the MIC-1 dendrite. It results from the crystallization of Ge between the Al grains but without the usual Al migration and without affecting the nature of the original top a-Ge layer. MIC-1 could be associated to defects in the interlayer, leading to Ge nucleation away from the Al-silica interface.

Interplay of MIC-1 and MIC-2/ALILE from EBSD and in-situ observations

From the in-situ observations (section 2.4 Figure 2.27 p. 40), the MIC-2/ALILE seems to progress regularly, starting from a single point and sweeping through large areas. shows how such an MIC-2/ALILE crystallization front progresses, starting from the left of the image and around a pre-existing Ge dendrite (created by MIC-1). In these conditions, we could expect the formation of a large single crystalline Ge platelet, all around the MIC-1 dendrite.

(a) (b) (c) (d) (e) (f) (g) (h) (i) (l) 20 μm (k) {101} {100} {111} substrate c-Ge (k)
In-plane EBSD Yet, in all the EBSD images, we never observe a single large grain around a dendrite, but several Ge grains with random in-plane orientation. This indicates that new crystals nucleate via MIC-2/ALILE as the crystallization front sweeps around the MIC1 dendrite. The schematic views in Figure 2.44 (a-e) summarize how such new nucleations may occur.

This observation complements the first in-situ images, in which the initial MIC-2/ALILE seed is always seen to nucleate close to the edge of a MIC-1 dendrite (see section 2.4.2.1 Figure 2.27 p. 40). The MIC-1 dendrite edge is thus always a likely nucleation location for the MIC-2 whether it is the primary nucleation as a starting point for the MIC-2 or a secondary nucleation generated when the ALILE crystallization front reaches a MIC-1 dendrite.

In practice, the high probability of MIC-2/ALILE nucleation at the MIC-1 dendrite edge causes all dendrites to be surrounded by several large Ge grains with random in-plane orientation, whether the dendrite is the initial nucleation center, or it is simply reached by an existing crystallization front.

The MIC-1 dendrite edge is a likely nucleation location for the ALILE/MIC-2 process, which leads to the formation of randomly in-plane oriented grains around each pre-existing MIC-1 dendrite.

Conclusions

Using EBSD and XRD, we have verified that both the MIC-2/ALILE and the MIC-1 are Metal Induced Crystallization processes which transform amorphous Ge into Ge crystallites.

In the case of the standard MIC-2/ALILE process, we confirm the literature results: we observe flat and large grains (several tens of microns wide) which exhibit a strong (111) texture. This is the standard orientation of the crystallites when grown on a silica surface with thin layers (Al layer less than 100 nm thick) [Toko, 2012;Kurosawa, 2012]. We have also confirmed that, when the layers exchange position, Al recrystallize in larger randomly oriented crystalline grains.

Given this new information about the crystallinity of Ge and Al, we can complete the preliminary results of section 2.5 (Figure 2.37 p. 52) to propose a complete description of the ALILE and MIC-1 mechanisms (Figure 2.45). ).

c-Al

AlGeO x a-Ge SiO 2 Al c-Ge ALILE / MIC-2 (b) MIC-1 a-Ge SiO 2 Al c-Ge Al (a)
Figure 2.45 -Schematics of the new MIC-1 process, which creates a crystalline Ge layer below the a-Ge layer [START_REF] Pelati | In Situ Optical Monitoring of New Pathways in the Metal-Induced Crystallization of Amorphous Ge[END_REF], compared to the standard MIC-2/ALILE process (b) which proceeds by the inversion of the Ge and the Al layers [Nast, 2000b].

Without high-resolution in situ optical monitoring, it is easy to mistake the MIC-1 region as a poorly oriented area created by a single ALILE process or by consecutive ALILE processes. In situ optical observations (section 2.4) combined with ex-situ analyses (sections 2.5 and 2.6) clearly demonstrate that both hypotheses are invalid, at least in our experimental conditions.

From the practical point of view, the MIC-1 phenomenon is detrimental to the quality of the layers obtained and its spatial extension need to be kept as small as possible.

Controlling size and position and with optical lithography

Should the lateral size of a single Ge grain match that of its substrate, we would obtain a Ge monocrystalline substrate on an amorphous support. Instead of trying to expand the grain size to unrealistic dimensions, we try here to adapt the size of the substrate using optical lithography to create domain of limited lateral extension, which can crystallizes independently.

This approach was already pursued by Cohin et al. on Si-MIC for the growth of III-V nanowires [START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF]. If the lateral dimension of each domain allows the nucleation and the growth of only a single grain, the final result is a regular array of independent monocrystals, of identical size and characteristics. Using well separated domains, we also avoid the formation of grain boundaries between adjacent grains. The use of micrometer-size domains could open the way to the fabrication of a single device on each of these micro-substrates.

We investigate in this section the process of MIC of Ge with Al, using optical lithography to obtain separate Ge monocrystalline domains. In particular, we address the question of the impact of the patterning on the previously observed MIC-1 and MIC-2/ALILE processes and how these affect the final morphology and crystal texture.

Lithography process

A first way to pattern the Al/Ge layers stack is to use the standard lift-off process. A first way to pattern the Al/Ge layer stacks is to use the standard lift-off process. The schematics of the patterned MIC fabrication is shown in Figure 2.46. The resist (AZ5214) is spinned on the sample (30 s at 4000 rotation/min). The substrate is either a thermally oxidized Si wafer (SiO 2 ∼40 nm) or a fused silica wafer. A soft bake for 1 min at 125 °C is used to drive off the excess of solvents and prepare the resist (Figure 2.46(b) (a)). Then, the sample is exposed for 5 s to UV light with the contact mask, baked for 1 min at 125 °C to reverse the exposed resist and exposed again to UV light for 30 s without the mask and finally developed (45 s in developer AZ826). After resist development (Figure 2.46(b)) and before material deposition, the sample is treated with a 1 min oxygen plasma to remove resist residuals in the holes.

We then deposit by e-beam evaporation the MIC structure, with the previously optimized stack of Ge/Al/AlGeO x /Ge layers of 1/20/1.5/20 nm thickness (interlayer oxidation for 9 min in air); see section (Figure 2.46(c)). The material between the domains is removed by lift-off in acetone using ultrasound sonication at room temperature (Figure 2.46 (d)). The sample is annealed at low temperature (250-300 °C) in order to crystallize the amorphous Ge (Figure 2.46 (e)). After the annealing, the Al layer is wet etched to expose the crystalline Ge patches(Figure 2.46 (f)).

Impact of the patterning on the MIC-1 process

MIC-1 and the size of the patterns

Description of the experiment In order to obtain monocrystalline patches, we would like to reduce the amount of MIC-1 dendrite nucleation, ideally to zero, and to have a single MIC-2/ALILE nucleation. We first investigate the effect of using different pattern sizes and their consequence on nucleation of the MIC-1 dendrite. Figure 2.47 shows Nomarski optical images of a sample composed of four different squares with respective widths of 40, 90, 175 and 275 µm.

Results

In all the squares, we observe the nucleation of several MIC-1 dendrites. These structures can be seen directly by optical microscopy (Figure 2.47 (a)), due to the slight contrast (light grey) or as a bright area in the SEM image (Figure 2.47 (b)), due to the associated increase of surface roughness.

Discussion

The spatial distribution of the MIC-1 dendrites is not uniform. Many dendrites are located at the edges of the pattern and relatively few are found inside the square area.

We thus compute two different densities of the Ge dendrites, separating the dendrites which nucleate inside the pattern elements and those which start from the edges. Figure 2.48 (a) shows the surface density of Ge dendrites nucleating inside the pattern. This is approximately a constant, d center = 1.2 • 10 4 µm -2 . For the dendrites at the edges, we compute a linear dendrite density (Figure 2.48 (b)) simply by dividing the number of such dendrites by the perimeter of the square. This also assumes an approximately constant value d edge = 0.04 µm -1 , i.e. one dendrite every 25 µm. Using these values, the total number of Ge dendrites is found to be:

(2.1)

d tot = A • d center + 4 √ A • d edge A
The linear edge density results in an increase of the total density of Ge dendrites for the small patterns (Figure 2.49). This simple model (blue dashed line), gives a good approximation of the measured total density of MIC-1 dendrites for each pattern (black squares). The edges of each patterned area are preferential nucleation sites for the MIC-1 dendrites. As a result, the total density of dendrites increases when the size of the pattern decreases.

Small patterns

Description of the experiment From the previous section, we estimate that by reducing the perimeter of a pattern below 25 µm (6.25 µm edge size for a square pattern), some patterns should have no MIC-1 Ge dendrite at all. This could permit the growth of pure MIC-2/ALILE grains, with no defective dendrite.

Results

We fabricate small squares of 2.7 µm side for which we expect a single Ge dendrites every 4-5 patterns. Figure 2.50 shows the corresponding square patterns observed in situ during the annealing. We clearly distinguish the onset of the formation of the MIC-1 dendrite between 25 min (Figure 2.50 (a)) and 1 h 30 min (Figure 2.50 (b)) annealing. Although each pattern (11 µm) is much smaller than the previously determined threshold value, each pattern presents a single MIC-1 dendrite (red arrow in Figure 2.50).

Conclusion

We have not managed to avoid the formation of MIC-1 Ge dendrites, even by reducing the pattern size well below the threshold value estimated from larger patterns. This might result from defective zones created by the lithography process which act as a preferential nucleation locus for the MIC-1 dendrites.

We cannot avoid the formation of the defective MIC-1 dendrites even on small patterns (2.7 µm), probably due to defects created by our lithographic process.

Impact of patterning on the MIC-2/ALILE

We now investigate if the MIC-2/ALILE is impacted by the patterning and the presence of the MIC-1 dendrites, first on a single large patch and then on an array of small patches.

MIC-2/ALILE on a large domain

Description of the experiment We look more in detail how the MIC-2/ALILE is impacted by the lithography process when using large rectangular patterns, 100 µm wide and 660 µm long.

Discussion

We observe comparable kinetics and spatial characteristics between the MIC-2/ALILE in a continuous layer and on large patches. The main difference is that the patches present more MIC-1 dendrites, which act as nucleation loci for the MIC-2/ALILE. This allows an easier nucleation, but it also reduces the area available for the ALILE and replaces it with poorly textured crystallites.

The MIC-2/ALILE process presents similar characteristics in continuous and patterned layers. The main difference is the increased number of MIC-1 dendrites in the patterned samples, which reduce the area available for MIC-2.

MIC-2/ALILE kinetics in patterned sample with a small pitch

Description of the experiment From the work of Cohin et al. [START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF] we expect that the pattern size and the etching damage process could have some effects on the crystallization kinetics, which may require an adjustment of the annealing temperature.

We now study the crystallization statistics of the small square patterns (2.7 µm large), previously presented in the section 2.7.2.2 (page 68). In a single in-situ image, we observe about 1700 individual patches, which allows us to quantify the nucleation statistics. We note that each individual patch already presents a single MIC-1 Ge dendrites (see section 2.7.2.2 p. 68) after 1 h at 250 °C, before any MIC-2/ALILE crystallization occurs. If the sample is kept at this temperature for another 11 h, we observe no further evolution (see insert in Figure 2.53 at t=6 h). Increasing the temperature from 250 °C to 269 °C induces the progressive nucleation by MIC-2/ALILE process in each pattern.

Discussion Similarly to Cohin et al. [START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF], if we assume that N (t) is the normalized count of crystallized patterns (MIC-2/ALILE) at time t and α the pattern crystallization rate, we find that: (2.2)

dN (t) = α(1 -N (t))dt
This simply means that the rate of increase of the number of crystallized patches is proportional to the number of remaining non-crystallized patches and to the nucleation rate. This yields:

(2.3)

N (t) = 1 -A 0 e -α(t-t 0 )
where A 0 is an integration constant. We find that A 0 = 1 because at t = 0, no patch has crystallized. From the fit (red line in Figure 2.53) we obtain a nucleation rate α = 7.86 • 10 -5 s -1 or τ = α -1 = 3.53 h. The fit is good in the final part of the experiment (t=17-38 h) but fails in the initial 6 h. The whole curve can be fitted by a sigmoidal function (green line in Figure 2.53). However, we do not have an explanation for the rather slow start of the nucleation. Overall, the crystallization rate is quite low and the "half life" of crystallization for this pattern is

t 1/2 = ln(2) • τ = 2.44 h.
As each crystallized patch requires a separate MIC-2/ALILE nucleation, the annealing duration for the complete inversion of the array (t>30 h) is much longer than for a large patch of equal surface (see Figure 2.51, p. 70) or for a continuous layer (see Figure 2.31 p. 44 in section 2.4.3).

The complete crystallization of array of small patches requires very long annealing times due to the low probability of MIC-2/ALILE nucleation in each pattern.

Crystalline texture of the patterned samples

Description of the experiment In the previous sections, we have investigated the basic morphology and the kinetics of the MIC process using lithography to limit the size of the grains. We study here if the lithography process has an impact on the crystal texture of the final Ge layer. Previous work by Cohin et al. [START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF] has shown that patterned layers could also adopt a (111) texture.

. Using EBSD, we characterize the crystal texture of the Ge grains formed in large (660×100 µm) and small (2.7×2.7 µm) patches, respectively studied in sections 2.7.3.1 and 2.7.3.2.

Results

Figure 2.54 shows the bottom of the large patch (660 × 100 µm) already studied in section 2.7.3.1, first observed in situ during the annealing (a-e) and then ex situ by optical microscopy (f) and by EBSD (g-h). The in-situ optical image, Figure 2.54 (a-e) shows that only one MIC-2/ALILE nucleation occurs in the selected part of the patch. Figure 2.54 (f) is a Nomarski optical image of the same area after standard wet etching of the Al layer (H 3 PO 4 ). We observe the expected double layer of Ge at the position of the MIC-1 dendrites (see Figure 2.54 (e)).

The corresponding EBSD map shows that the crystal texture in the vertical direction (Figure 2.54 (g)) is a mixture of ( 100) and ( 111) oriented small grains (<1 µm), despite the fact that they result apparently from a single MIC-2/ALILE nucleation. The in-plane orientation of the grains is random (Figure 2.54 (h)). In both EBSD maps, the MIC-1 dendrites appear dark (no EBSD signal) due to the amorphous top Ge layer (see Figure 2.43 in section 2.6.2.3). A similar MIC structure of a continuous layer (not processed by lithography) shows a clear (111) texture and large grains (>10 microns wide) (Figure 2.54 (i-j)).

Figure 2.55 shows the SEM image and corresponding out-of-plane and in-plane EBSD maps of a single square patch (2.7×2.7 µm). As in the larger patches (Figure 2.54), we observe a random crystal orientation, i.e. no preferential (111)-texture, and very small grains (1 µm and less).

Discussion

The strong (111) crystalline texture of the continuous layer is completely lost when the lithography is used to limit the area of the initial bilayer, whether we use large or small patterns. This important change occurs although the in-situ optical images are not different. This shows that a single MIC-2/ALILE nucleation can actually produce randomly oriented grains.

(f) Nomarski optical image after Al wet etching. Comparison of EBSD analysis of the Ge patch (g,i) and a reference non-patterned sample (see section 2.6.2.1 p. 57) (h-l). EBSD in the vertical direction (g-i) and in-plane direction (h-l)), i.e. parallel to the substrate.

A likely culprit for the loss of the (111) texture is a change of the silica-Al interface, which is known to determine the final crystal orientation [Toko, 2014b]. Our liftoff process (see section 2.7.1, Figure 2.46), p. 65) relies on the deposition of a resist, which is later developed and exposed for 1 min to an oxygen plasma. Despite these cleaning steps, it is possible that contaminants could still be present on the silica and affect the crystalline texture. This can be tested by devising other lithography schemes, which do not affect the Al-silica interface. The lithographic process based on lift-off results in the complete loss of the (111) texture. We only obtain randomly oriented crystals of a size smaller than those obtained using continuous layers..

Alternative lithographic process

To improve the crystal quality of the final Ge patches obtained in patterned substrates, two problems need to be addressed:

• The reduction of the formation of MIC-1 dendrites on the edges of the patches (see section 2.7.2)

• The conservation of the (111) crystal texture (see section 2.7.4).

Edge geometry and MIC-1 dendrites

Description of the experiment To investigate the cause of the increased density of MIC-1 dendrites along the edges of the pattern, we prepare two different edge geometries using silicon wafer pieces to mask selective parts of the sample during the e-beam deposition. In the first one, we simply mask the same area during Al and Ge depositions (Figure 2.56 (a)). In the second sample, we move the mask between Al and Ge deposition to obtain a staircase geometry, in which the Al layer extends further than the Ge layer (Figure 2.56 (b)). The Al layer include the Ge-underlayer and the Ge(AlOx)-interlayer.

Results

In Figure 2.56, we compare the Nomarski images of the two geometries of a MIC-Ge stack after 4 h of annealing at 250 °C. In Figure 2.56 (a), showing the full bilayer edge, we observe the usual enhanced nucleation of MIC-1 dendrites. At the staircase edge (Figure 2.56 (b)), the density of MIC-1 dendrites is much lower.

If only the Ge layer is patterned, we observe the same dendrite density on the edge of the patch as inside.

Finally, we turn to the edge of the Al layer (Figure 2.56 (c)). There, the Al layer is covered by the Ge(AlOx) interlayer. Despite the absence of the main Ge layer, numerous small dendrites nucleate right at the edge of the Al layer.

Discussion

If the patterning is limited to the top Ge layer (Figure 2.56 (b)), we recover the standard nucleation rate for the MIC-1 dendrites. This indicates that the edge-enhanced MIC-1 dendrite nucleation observed section 2.7.2.1 is caused by the discontinuity of the Al layer. The nucleation of numerous dendrites at the edge of the sole Al layer (Figure 2.56 (c)) confirms this hypothesis. This proves that MIC-1 dendrites are associated to defects present in the Al layer (or very close to it), either in the Ge underlayer or in the Ge(AlOx) interlayer.

We find that the patterning of the Al layer is responsible for the enhanced nucleation of the MIC-1 dendrites at the edges of the patches.

Revised patterning process for the etching of the top Ge layer

Following the previous findings, we switch to a positive lithography in which we only pattern the top Ge layer. First, we deposit the Al/Ge layers stack on the silica (Figure 2.57 (a)), then we spin the resist (AZ5214) and anneal the sample (soft-bake) for 1 min at 125 °C (Figure 2.57). We then insolate the sample and we develop the apertures in the resist. After that, we wet etch the exposed Ge in H 2 O 2 for 180 s to create the pattern in the top layer. Finally the resist is stripped using acetone (5 min, with ultrasound). At this stage, it is also possible to etch the exposed Al layer in H 3 PO 4 before removing the resist. The sample is then annealed, as before, to induce the MIC of amorphous Ge. Then, Al is wet etched using the standard H 3 PO 4 solution.

We now investigate the effect of this new lithographic process on the MIC-1 and the MIC-2/ALILE morphologies.

Effect of the revised patterning process on MIC-1 and MIC-2

We now repeat the annealing experiment using the sample fabricated with the revised patterning process.

Results Figure 2.58 shows steps in the annealing of a large patch (100x600 µm).

The sample is annealed 260 °C. The stack deposited Ge/Al/GeOx/Ge 1/20/1.1/22 nm thick, the interlayer being oxidized for 10 min.

We observe in Figure 2.58 (b,c) (white arrows) that the MIC-1 nucleation occurs in the middle of the pattern. Some nucleation is also visible outside the patch (purple arrows in Figure 2.58 (c)), possibly due to the Ge underlayer or to diffusion of Ge from the patch (see orange arrow in Figure 2.58 (d)). The MIC-2/ALILE starts as usual from one of the available MIC-1 dendrites (those which are not on the edge (red arrow in Figure 2.58). The MIC-2/ALILE continues until the whole patch is crystallized (Figure 2.58 (e-h)), as in the case of the previous patterned samples (Figure 2.51).

In the case of square patches 40 µm wide, we observe only one dendrite per patch, which is located inside the patterned area (Figure 2.59 (d), instead of several dendrites, all at the edge of the patch (Figure 2.59 (c), for the previous lift-off procedure.

Discussion

The revised etching procedure leads to a significant reduction of the density of MIC-1 dendrites. It completely eliminates the edge-related nucleation (see section 2.7.2). This confirms the preliminary results of section 2.7.5.1, obtained by masking the sample during the deposition.

A drawback of this new procedure is the parasitic nucleation of new MIC-1 den-drites, outside the patterned area. This creates additional Ge crystals, outside the masked area, after etching of the Al layer. This shows that the Ge material used in MIC-1 can originate from other sources that the main Ge layer, which may be either the Ge underlayer or the Ge(AlOx) interlayer.

The crystallization of patterns obtained by the revised etching procedure show much reduced density of MIC-1 dendrite, due the complete elimination of the edge-enhanced nucleation process. In turn, this should lead to a much crystal improved texture of the final Ge layer.

Effect of the revised patterning process on the crystal orientation

Description of the experiment We now characterize the crystal texture of the patterned samples made using the revised lithography process (limited to the Ge top layer) (section 2.7.5.2 p. 76). Since no resist is deposited on the silica surface, we expect minimal deviation from the preferential (111) texture obtained after the optimization on continuous layers (section 2.6.2.1). GIXRD for the ( 111) and ( 100) crystal texture of the samples (#1) and (#2) reported in the Figure 2.39 We then characterize the crystal texture using GIXRD. The corresponding diffractrograms are shown in Figure 2.60 (c). The raw detector intensity is lower for the patterned sample, despite similar substrate sizes, which is expected as only a small fraction of the total surface is patterned with the Ge crystals. The relative diffracted intensities are similar, with a dominance of the 220, 422 and 440 peaks, which is indicative of a strong (111) crystal texture (see section 2.6.1 p. 2.6.1). The quantification of the crystal texture using Lotgering factors for the (111) and (100) orientations is shown in Table 2.7. The Lotgering factor for the patterned layer is higher for the (111) texture than that of the reference continuous layers.

Results

Lotgering factors computed from GIXRD diffractrograms are not very selective between the (111) and the (100) textures because the later have many common diffracted intensities, see Appendix A. Still, the higher values for the (111) texture indicate a larger probability of (111)-oriented grains.

Discussion Contrary to the previous lithography procedure using lift off (section 2.7.4, p. 73), which leads to a complete loss of the (111) texture, the revised patterning procedure maintain a high (111) crystal texture. This confirms that the MIC process is extremely sensitive to minor contamination of the starting silica surface. The standard cleaning procedures based on chemical and plasma cleaning are thus not suitable: they either fail to remove trace amounts of organic compounds from the resist or they cause surface damage, which leads to the loss of the (111) crystal texture in the final Ge platelets.

Samples patterned using the revised lithography procedure conserve the strong (111) crystal texture found in the continuous layers.

Conclusion

The idea of obtaining monocrystals of defined size and position by patterning the MIC layers is challenging. In a first attempt using a lift-off technique to pattern the samples, we identified two problems:

1. The extra-nucleation of numerous MIC-1 dendrites on the edges of each patch, which lowers the overall crystal quality by creating zone of dense, poorly oriented, small grains.

2. The complete loss of the (111) texture for the large Ge grains crystallized by the MIC-2/ALILE process, despite the robust results obtained on continuous layers.

We thus revised our etching procedure to etch only the top Ge layer and to avoid tampering the silica surface before Al is deposited. This revised etching process leads to a dramatic decrease in the MIC-1 dendrites and it has restores the preferential (111) orientation of the grains.

The final shape of the grain is not regular and it only loosely matches that of the pattern features. Improving surface coverage and grain size would require further studies.

EBSD could not be performed on the final patterned sample layer due to time constraints and we still need to assess if the obtained patches are monocrystalline. However, the presence of MIC-1 dendrites inside the patch most likely results in the growth of several sub-grains with random in-plane orientation (see section 2.44 p. 62).

Conclusion on MIC fabrication

Metal Induced Crystallization (MIC) is a technique which allows amorphous materials in contact with a metal layer to crystallize at very low temperature. We focused on the fabrication of mono-crystalline and strongly textured Ge platelets. In the present case, we used Al to catalyze the crystallization, with typical annealing temperatures of 250-300 °C, well below the Al-Ge euctectic (421 °C).

The literature reports several interesting examples of the Al-induced crystallization of rather thick layers (50 nm), resulting in a high (111)-texture, but also highlights various problems encountered with thinner layers. In this work, we showed that thin Ge layers (20 nm) can be crystallized using the classical layer exchange process (ALILE) provided several parameters are adjusted. In addition to the thickness of the main semiconductor and metal layers, the nature of the interlayer deposited between the Ge and Al layers plays a critical role in regulating the diffusion of Ge atoms. To obtain a highly textured Ge layer, the Ge atoms must nucleate at the Al-silica interface and not at the Al grain boundaries. We showed that the standard AlO x interlayer must be replaced by a Ge(AlO x ) interlayer to restore the kinetics of the ALILE process. We also need to add a Ge underlayer, between the Al and the silica, to promote the ALILE process rather than a new parasitic crystallization mechanism that we discovered and named MIC-1.

Beside the successful crystallization of thin Ge layers (20 nm), less than half as thick as those reported in the literature, one of the main features of this work is the development of a dedicated in-situ microscope to image and measure the progress of the crystallization during annealing. This instrument has proven invaluable in identifying two distinct crystallization mechanisms, the standard ALILE (here called MIC-2) and the new MIC-1 parasitic process. By combining various ex-situ characterization techniques (TEM, EDX, EBSD), we are able to fully determine the structure created in each crystallization pathway: a bilayer of crystalline Ge and Al grains for ALILE, with a strong (111) texture, and Ge crystals covered by amorphous Ge for MIC-1, with mixed (100) and (111) orientations.

Finally, we investigated the possibility of crystallizing regular arrays of patches of controlled size and position, with the idea to reduce, or even eliminate, the boundaries between different grains of the MIC. In a first attempt, the etching procedure lead to an enhanced nucleation of MIC-1 dendrites at the edge of each patches, as well as a total loss of the (111) texture. After identifying the possible causes of these problems, we implemented a revised procedure, which eliminates MIC-1 nucleation on the edge of the patches and restores the (111) texture of the final Ge platelets. Despite these significant improvements, we are still not sure if each patch consists of a single Ge crystal.

Starting from a bare silica surface, we succeeded in creating arrays of regular patches of crystalline (111)-oriented Ge, several tens of microns wide. As regards subsequent III-V epitaxy, the present surface is a crude approximation of a bulk Ge monocrystalline (111) wafer, but it only uses 1/10000 of Ge material, namely a thickness of 20 nm compared to 200 µm. Our attempt to fabricate fully monocrystalline and well-separated micro-substrates met some successes but still calls for improvement.

Despite this limitation, we will investigate in the next chapter how these MIC-Ge pseudo-substrates compared to standard Ge (111) commercial wafers for the epitaxy of GaAs material using molecular beam epitaxy.

Chapter 3

GaAs & Ge epitaxy by MBE on MIC-Ge and reference substrates

In this chapter we investigate the epitaxy of GaAs on standard GaAs(111)B, Ge wafers and on MIC-Ge pseudo-substrates. Several problems need to be solved before we tackle the epitaxy of GaAs on patterned MIC-Ge.

We first investigate if the Ge layer fabricated by MIC can be improved, in terms of roughness, texture and crystalline structure. Toko et al. showed that it is possible to obtain a 500 nm thick epitaxial Ge layer on a MIC-Ge layer. However, the crystalline quality of this layer depends a lot on that of the MIC-Ge seed layer [Toko, 2015].

Secondly, we investigate the epitaxy of GaAs on MIC-Ge pseudo-substrates by MBE. This is done in three steps:

1. Homo-epitaxy of GaAs on GaAs (111)B monocrystalline wafers, in order to investigate the growth parameters of GaAs layers in the (111) orientation;

2. Hetero-epitaxy on Ge (111) wafers, in order to study the growth of polar GaAs, on non-polar Ge (111) surfaces.

3. Selective GaAs growth on MIC-Ge pseudo-substrates on silica, to limit the GaAs growth to the oriented crystalline Ge patches so that we avoid random GaAs nucleation on the amorphous silica surface.

Improving the MIC Ge surface by deposition of an epitaxial Ge layer

In this section, we attempt to improve the roughness and crystal quality of the continuous or patterned MIC-Ge pseudo-substrates by depositing an homoepitaxial Ge layer using MBE. In the case of patterned substrates, the growth should be selective, i.e. occur on the Ge patches but not on silica.

We preliminary characterize by XPS the thermal deoxidation of the Ge. Then, we deposit the Ge on MIC-Ge layer with low Ge coverage of the silica. We focus on selectivity and on the final roughness. The Ge crystal structure is diamond cubic, similar to that of Si. The Ge lattice parameter is 5.658 Å1 . Figure 3.1 and Figure 3.2 shows the different views of the conventional unit cell of cubic Ge.

Ge crystal structure

The Si-Ge growth chamber

The MBE set-up

The epitaxy of Ge was performed in a chamber dedicated to Si and Ge growth (Figure 3.3), connected to the main chamber of a Compact21 RIBER MBE fitted for III-V epitaxy. The chamber is composed by: the main viewport, the viewport to see the Si crucible, the Medium High Temperature (MHT) Ge cell, the manipulator (external and internal view), the gauge pressure sensor, the RHEED viewport with phosphorous screen, the electron gun for Si evaporation, the pyrometer, the quartz sensor, the quartz oscillator, and a free cell port. On the manipulator we see: the rotating parts, the oven and the molybdenum sample holder.

Ge flux calibration

To perform growths in controlled conditions we calibrate the Ge growth rate versus the Ge cell temperature. In the Si-Ge chamber, there is a single pressure gauge to measure the pressure in chamber but none dedicated to the measurement of the cells flux. However, we can relate the chamber pressure with the flux impinging on the sample using the quartz measurements or ex-situ SEM observations. Figure 3.4 (c) shows the chamber pressure recorded as a function of the cell temperature.

To correlate these parameters with the growth rate, we fabricate a sample with several Ge layers grown at different cell temperatures on a Si (100) substrate. These We verify the Ge growth rate on Ge(100) substrate using reflection high energy electron diffraction (RHEED) oscillations (see Figure 3.5). For this measurement, a low substrate temperature is necessary [Daniluk, 1996], about 300 °C. From the RHEED, we obtain a growth rate 0.289 µ/h at a cell temperature of 1300°C in agreement with the previous SEM calibration.

(a) (b)

Ge @ 1300 °C → 0,289 µm/h

Ge surface preparation -deoxidation and carbon removal

Description of the experiment

We verify here if the Ge layer made by MIC can be thermally deoxidized. To this purpose, we heat the sample and we characterize it by XPS without breaking the vacuum. We performed a series of annealing experiments at 450 °C, 650 °C, 700 °C, 750 °C and 800 °C on the same sample. At each annealing step, the sample is moved from the XPS chamber to the deposition chamber and vice-versa, (Figure 3.6). We focus on the 3d and 2p XPS peaks of Ge which energy is 29-30.2 eV [Prabhakaran, 1995;Molle, 2006a;[START_REF] Oh | Thermal desorption of Ge native oxides and the loss of Ge from the surface[END_REF]Hovis, 1999] and 1217. 3-1217.6 eV [Prabhakaran, 1995;[START_REF] Oh | Thermal desorption of Ge native oxides and the loss of Ge from the surface[END_REF]Hovis, 1999] respectively. We also consider at the O 1s and C 1s peaks respectively at 530.8-532.4 [Prabhakaran, 1995;[START_REF] Oh | Thermal desorption of Ge native oxides and the loss of Ge from the surface[END_REF] and 285.8 eV [Hovis, 1999]. Ge can form four type of oxides, each one with a different shift of the Ge 3d peak [Molle, 2006b;[START_REF] Oh | Thermal desorption of Ge native oxides and the loss of Ge from the surface[END_REF]:

Ge 2 O (Ge 1+
) with a shift of 0.7-0.8 eV; GeO (Ge 2+ ) with a shift of 1.7-1.8 eV; Ge 2 O 3 (Ge 13+ ) with a shift of 2.6-2.8 eV and GeO 2 (Ge 1+ ) with a shift of 3.4-3.5 eV. The Ge 2p peak is reported to shift from 1217.6 eV to 1219.3 eV for the GeO oxide and to 1220.6 for the GeO 2 oxide [Prabhakaran, 1995]. Results Figure 3.7 (a) shows clearly a double peak of the Ge 3d before annealing, which can be attributed to Ge and GeO 2 . After a first annealing at 450 °C, the GeO 2 oxide seems to disappear and the Ge peak is more intense. Then, the spectra show very little changes up to 750 °C. Above 800 °C, the Ge signal drops significantly.

25
A similar situation for the Ge 2p peak is visible in Figure 3.7 (b). Before annealing, we observe a peak from GeO 2 at 1221 eV and a peak from pure Ge at 1218 eV. After annealing at 450 °C, the oxide contribution almost disappears and it vanishes totally after annealing at 650 °C. The pure Ge peak increases to full height at 450 °C and does not change up to 750°C, and almost disappears after the annealing at 800 °C.

The oxygen peak in Figure 3.7 (c) is clearly visible before annealing and after the first annealing at 450 °C, but with reduced intensity. It continues to decreases up to 700 °C and then increase again at 800 °C.

The carbon peak intensity is strongly reduced after the first annealing at 450 °C, (Figure 3.7 (d)) and further reduced upon annealing at 700 °C and 750 °C.

Discussion

The observed shifts of the Ge 2p and Ge 3d peaks are due to the native oxide on the MIC-Ge patches and indicate that this oxide is mainly GeO 2 , as expected from the literature [Deegan, 1998;[START_REF] Oh | Thermal desorption of Ge native oxides and the loss of Ge from the surface[END_REF]. Ge is reported to completely deoxidize at 500 °C [Prabhakarana, 1994;Prabhakaran, 1995;Molle, 2006b] in vacuum. Although the GeO peak starts to disappear at 400 °C [Prabhakaran, 1995] and the main reduction occurs at 450 °C [Molle, 2006a]. In our case we don't see any significant oxide desorption after the first annealing at 450 °C which is coherent with the literature results.

The GeO 2 peak is reported to vanish at 475 °C [Molle, 2006a] although the oxide should be stable up to 1170 °C [START_REF] Oh | Thermal desorption of Ge native oxides and the loss of Ge from the surface[END_REF]. This is explained by the reaction of the germanium dioxide with Ge to produce germanium oxide GeO (GeO 2 + Ge → 2GeO). The chemical reaction requires Ge material, so that no desorption of GeO 2 on SiO 2 occurs even if annealed at 550 °C [Wang, 2010].

Most of the literature reports only on Ge(100), but similar Ge oxide desorption kinetics are observed for Ge(100) and Ge(111) surfaces [Deegan, 1998;Prabhakaran, 1995].

Ge can also chemisorb carbon and form (Ge,C) compounds [Amy, 2006]. From the C 1s peak variation, we see that the main carbon contamination is desorbed after the second annealing at 650 °C.

After the annealing at 800 °C the two Ge peaks are greatly reduced and the O 1s is raised. This is explained by the evaporation of most of the ∼20 nm MIC-Ge which exposes the underlying silica surface to the XPS beam.

• Most of the germanium oxide desorbs at 450 °C and completely desorbs at 650 °C.

A small oxygen signal persist due to the non complete coverage of the silica by the MIC-Ge.

• MIC-Ge layers are stable up to 750 °C and starts desorbing at 800 °C.

• The carbon chemisorbed on the germanium surface is efficiently desorbed by an annealing at 650 °C.

Short review on Ge epitaxy

We now examine the recent literature to find appropriate growth conditions for the epitaxy of Ge.

Ge homoepitaxy in literature

On MIC-Ge Toko et al. [Toko, 2015]reported the homoepitaxial growth of Ge, highly (111) textured, using a substrate temperature at 500 °C.

On Ge (100) layer, 6°off, Masuda et al. [Masuda, 2016] studied the growth conditions of As-doped Ge layers. Up to 500 °C the growth shows a step-bunching behavior with decreasing surface roughness as the substrate temperature increases. In the range of 550-600 °C the growth proceeds by step-flow. An additional annealing at 600-640°C after growth can also improve the surface of the epitaxial Ge layer [Masuda, 2016;Sieg, 1998a].

Conditions for selective growth on silica

In the latter stage of this work, we plan to use patterned MIC-Ge layers on silica substrates. We investigate here the homoepitaxy of Ge on individual MIC-Ge platelets crystallized on a silica layer (Figure 3.8). In order to selectively grow Ge on the existing Ge patches and not on the silica substrate, we need to adjust the growth parameters. It is shown that for a thick thermal silica layer (> 5 nm), the impinging Ge does not react with the silica to form volatile GeO and SiO [Li, 2006b;Li, 2006a]. As we use here 40 nm thick silica layers, the Ge selectivity does not rely on the formation and evaporation of Ge oxides but on the fast desorption of Ge adatoms on the silica surface.

To avoid the formation of stable Ge nuclei on SiO 2 , we have to consider the impiging flux and desorbing Ge fluxes on the silica. The first one depends on the Ge cell temperature and the second depends exponentially on the substrate temperature [Li, 2006a;Li, 2006b]. Figure 3.9 shows the threshold where the desorption flux exactly counterbalance the impinging Ge flux for different substrate temperature. A higher impinging flux leads to formation of stable Ge nuclei whereas for lower fluxes, the Ge adatoms desorb before forming stable nuclei. [Li, 2006b] Therefore, we must use a high substrate temperature and a low Ge flux to selectively grow on the existing Ge patches and not on the surrounding silica surface. In the following, we choose a substrate temperature of 550 °C (pyrometer) and we set the Ge cell around 1200°C (growth rate of 0.12-0.16 ML/s). The blue point in Figure 3.9 represents our growth conditions.

Experimental results

Substrate cleaning

The MIC-Ge surface is first cleaned with an HBr solution. This solution dissolves all the Ge sub-oxides (GeO x , x< 2), with only a limited etching of Ge [Onsia, 2005]. HBr also provide a better passivation than HF for Ge surfaces [Jagannathan, 2006]. Prior the growth, the samples are annealed at 640-700 °C in a separate degassing chamber to obtain a complete deoxidation [Sieg, 1998b;Kawai, 1993].

Growth morphology and selectivity

Description of the experiment We first characterize by SEM a Ge layer grown on a continuous MIC-Ge substrate to evaluate the surface roughness and the growth selectivity. The MIC consists in well isolated Ge grains on the silica. We grow for 2 h 20 min using a Ge cell temperature of 1200 °C (about 350 nm at ∼0.15 ML/s), see Figure 3.10.

Then, we compare two patterned MIC-Ge substrates with large patches (side > 40 µm) grown at the same substrate temperature (550 °C) but using different Ge fluxes: Ge cell at 1200 °C and 1240 °C. Since the 1240 °C Ge flux is about twice that of 1200 °C, we reduce the growth time from 2 h 20 min to 1 h, to obtain similar final Ge thickness (Figure 3.11). We clearly observe Ge growth on the MIC-Ge platelets, in the form of grains of 1-3 µm large. On the silica, we observe small Ge crystallites about 0.1-0.3 µm wide.

Figure 3.11 shows SEM top views and tilted views of a large MIC-Ge patch. Here, the Ge growth was performed with low (a,c) and high (b,d) fluxes. The growth morphology is similar in both samples. The grains are slightly larger on the sample grown with the lower flux. The large Ge flux is associated to higher density of small crystallites on the silica surface.

Discussion Despite the chosen favorable growth conditions for the selective growth of Ge with respect to SiO 2 surface (substrate temperature > 500 °C and relatively low Ge flux), we observe that Ge crystallites nucleates on the silica. On the MIC-Ge patches, the homepitaxial Ge growth creates isolated grains (1-3 µm wide) instead of a continuous layer. This contrasts with the results reported for the epitaxy on Ge (100) wafers [Masuda, 2016;Sieg, 1998b] or on (111) MIC-Ge pseudo-substrates [Toko, 2015]. • The resulting high surface roughness is not appropriate for the subsequent III-V growth. The relatively poor selectivity is also problematic for the GaAs epitaxy, as all the Ge crystallites on the silica will also act as a seed for the GaAs crystals. Despite a low Ge flux and a high substrate temperature, the growth of Ge is only partially selective with respect to the silica surface.

• The growth of Ge on the MIC-Ge patches creates individual grains, which do not coalescence within the growth duration.

The later significantly increases the roughness of the surface and is a poor starting point for the subsequent GaAs epitaxy.

EBSD characterization

We perform EBSD on the grown Ge layer to find if the growth is epitaxial and if the (111)-texture is maintained. Result Figure 3.12 (a,b) shows the crystallographic orientation of the grains in the direction normal to the substrate (a) and in-plane (b). We observe a predominant (111) texture normal to the substrate (blue color). A minority of the grain is oriented in the <100> direction. (red color). All the grains have a random in-plane orientation.

Discussion

We observe that the partial (111) crystal texture of the MIC-Ge pseudo-substrate is transfered to the top Ge island obtained after the Ge growth by MBE. This supports the hypothesis that the growth of Ge occurs in epitaxy on the existing Ge islands, as previously reported by Toko et al. [Toko, 2015]. However, cross-section transmission electron microscopy of the interface would be required to fully demonstrate this hypothesis.

The Ge material grown by MBE on the top of the MIC-Ge patches show a partial (111)-texture, similarly to that of the MIC-Ge pseudo-substrates. This support the hypothesis of epitaxial growth.

Conclusion

We investigated the homoepitaxy of Ge on the MIC-Ge patches obtained by MIC on silica. We first characterized the thermal deoxidation of Ge and verified that it is complete at 650 °C.

Then, we perform the MBE growth of Ge on the MIC-Ge template. We observe that the Ge material grows in the form of individual islands rather than the expected continuous layer. Moreover, the selectivity is not perfect and some Ge crystallites also appear on the silica. The similar crystal texture of the grown material with that of the MIC-Ge pseudo-substrate is compatible with epitaxial growth.

Considering the high surface roughness and the partial selectivity of the Ge growth, we decide to avoid the growth of such Ge buffer layer and we investigate the direct growth of GaAs on the MIC-Ge pseudo-substrates.

Basics of GaAs growth by MBE

To obtain a GaAs solar cell on our MIC-Ge pseudo-substrates, we need to control the heteroepitaxy of GaAs on Ge in the <111> direction. However, the later is not the most commonly studied, compared to the <100> orientation.

To better understand the impact of the MIC-Ge substrates on the GaAs growth, we will first study the homoepitaxy on GaAs(111)B single crystal substrates and the growth of GaAs on Ge (111) monocrystalline wafers.

In this section we recall some basics on the GaAs crystal structure and then we describe the epitaxy setup and the calibration of the molecular fluxes. The crystalline structure of GaAs is zinc blende. As for the diamond cubic structure found in Ge, the zinc blende consists of two face centered cubic (fcc) sublattices. The two sublattices are shifted along the diagonal of the cube by a quarter of its length,

GaAs crystal structure

a 0 √ 3
4 . One sublattice is occupied by Ga atoms, the other by As atoms. The conventional cubic unit cell is shown in Figure 3.13. Different crystallographic orientation of atomic stacking are shown in Figure 3. 14 (a-c). Ge has the advantage to be closely lattice matched with GaAs (5.65325 Å for GaAs and 5.658 Å for Ge), and the thermal expansion coefficients of both materials are also close (5.73 • 10 -6 K -1 for GaAs and 5.9 • 10 -6 K -1 for Ge).

The two different atoms having different electronegativities, GaAs is a polar material. In particular, the top of the (111) surface is composed of either Ga atoms, (111)A, or As atoms, (111)B, see Figure 3.14. The growth conditions on these two (111) surfaces are different.

Molecular Beam Epitaxy chamber for III-V materials

GaAs is grown in a Molecular Beam Epitaxy (MBE) equipment. Ga and As atoms are evaporating from solid sources. The generated fluxes are directed toward the sample surface. chamber chamber: a cryogenic pump, a ionic pump and a cryopanel cooled with liquid nitrogen.

III-V flux calibration

In any quantitative study, it is important to calibrate the atomic fluxes impinging on the sample. We describe here the procedure for the As and Ga fluxes calibration.

We measure the fluxes of the evaporated materials with the pressure gauge and we measure the associated planar growth rate using RHEED. We first calibrate the Ga flux as a function of the cell temperature and the As flux as a function of the opening valve.

Thanks to the very small impinging angle of the electron beam (1-3 °), the RHEED only probes the topmost atomic layers. For crystalline layers, the main interaction between the beam and the surface is that of diffraction. For rough surfaces, we observe the classical dotty figure which matches the projection of the full reciprocal lattice on the RHEED screen, following the Ewald sphere construction. The diffraction of smooth layers reduces the diffraction pattern to vertical rods, where each one corresponds to one particular diffraction spot, spread in the normal direction. The later is essentially the Fourier transform of a 2D crystal in 3D reciprocal space.

The growth rate is measured from the oscillations of the intensity in selected regions of the diffraction pattern. These oscillations are related to the change in surface morphology during the layer by layer growth mode. When no adatom is present on the surface, the diffraction pattern is that of rods (2D diffraction). Then during the formation of a monolayer the surface is rougher (adatoms have not yet formed a full monolayer) and the diffraction pattern tends to coalesce into dots (3D diffraction). As a consequence, the diffracted intensity in selected region of the RHEED screen oscillates with the formation of each new monolayer. This enables to measure in-situ the exact atomic growth rate and to compute the atomic flux on the sample for each effusion cell temperature.

For Ga, we record the pressure gauge measurement, named beam equivalent pressure P BEP , as a function of the cell temperature. We observe an Arrhenius behavior with an exponential variation as a function of the inverse of the temperature. The relation is shown in equation 3.1.

(3.1) log(P BEP [T orr]) = a p + b p T [K]
Experimental values are shown as black points in Figure 3.17. It is shown also a linear fit of the experimental value by a red line. We then measure the growth rate in situ using RHEED, using standard GaAs (100) substrates, for different Ga cell temperatures (Figure 3.18). We limit the growth to the As-rich conditions (excess of As) so that the GaAs growth rate only depends on the impinging Ga flux.

We can now combine the measured beam equivalent pressure (BEP) with the measured growth rate for the same temperatures of the Ga cell. The relation between BEP and growth rate is linear (Figure 3.19):

We can now combine the measured pressure (BEP) with the measured growth rate for the same temperatures of the Ga cell. The relation between BEP and growth rate is linear (Fig. 3.19):

(3.2) g r [Å/s] = q + w • P BEP [T orr]
with q = -0.0464 and w = 1.088•10 -7 . Within the experimental error, we observe that the Ga BEP is strictly proportional to the (Ga-limited) GaAs growth rate. 0,00081 0,00082 0,00083 0,00084 0,00085 0,00086 0,00087 0,00088 Using the ML thickness (a 0 /2) and the Ga atom density (2/a 2 0 ) for (001) planes, we finally compute the relation between atomic flux of Ga atoms Φ Ga and the measured BEP value P BEP .

-2,0 -1,8 -1,6 -1,4 -1,2 -1,0 -0,8 -0,6 -0,4 -0,2 0,0 ln(r [Å/s]) Ga 1/T (K -
(3.3) Φ Ga = 4 • g r /a 3 0 = 2.41 • 10 7 • P BEP (Ga)
with Φ Ga in at. nm 2 •s and P BEP in Torr. We now keep the Ga flux constant (Ga cell at 935 °C, BEP ∼1.44•10 -7 Torr) and we vary the As valve opening, expressed here in percentage of the aperture range. The substrate temperature is at 580-585 °C (pyrometer). The temperature of the cracker of the As cell is always 600 °C, so that we evaporate only the As 4 molecules. Below a certain value of the As flux, the growth rate is limited by the As and it decreases with each further reduction of the As valve aperture (Fig. 3.20). The transition point between the constant growth rate (Ga-limited) and the variable growth rate (As-limited) with respect to the valve aperture is the point for which the As and Ga atomic fluxes are equal. GaAs growth rate (Å/s)

Ga limited As limited

Figure 3.20 -Planar GaAs (100) growth rate measured at 580 °C from RHEED oscillations as a function of the As valve aperture. The black vertical line indicates the valve aperture between the As limited regime (left) and the Ga limited regime (righ

We now plot in Figure 3.21 the As BEP measured with the pressure gauge as a function of the aperture of the As cell valve. The curve is not linear and can be fitted with a third order polynomial (red line).

We now combine the measured growth rate in the As limited regime (Figure 3.20) with the BEP measured for the same valve opening to obtain the variations shown in Figure 3.22.

Similarly to the Ga-limited regime, the GaAs growth rate in the As-limited regime is linear with the As 4 BEP but it is not strictly proportional. As the As and Ga atomic density are equal in GaAs crystals, we compute the As atomic flux from the GaAs growth rate in the same way than equation 3.3 to obtain:

(3.4) Φ As 7. GaAs growth rate (Å/s)

Homoepitaxy on a GaAs (111)B surface

In this section we investigate the homoepitaxial growth of GaAs on GaAs(111)B substrates. The growth is first characterized in situ using the RHEED. Then we evaluate the impact of the growth temperature and the As flux on the roughness. We first shortly review the results from the literature, with a focus on the RHEED reconstructions observed in the GaAs(111)B and the surface morphology as a function of the growth parameters. 111)B (from [Woolf, 1993a;Woolf, 1993b]). (b) Table of RHEED reconstructions as a function of the growth temperature (from [Woolf, 1993a]).

The RHEED reconstructions observed for the GaAs(111)B surface are not only function of the substrate temperature but also of the III:V ratio [Woolf, 1993a;Woolf, 1993b;Yang, 1992] (Figure 3.24 (a)).

Generally the (2×2) reconstruction appears at low substrate temperature and high As flux and the ( √ 19x √ 19) reconstruction at high substrate temperature or low As fluxes [Yang, 1992]. Different models has been proposed to explain the ( √ 19x √ 19) reconstruction, [Koga, 2010].

Another (1x1) regime appears at high temperatures (>630 °C), see Figure 3.24. 111)A (from [Woolf, 1993b]). (b) Table of RHEED reconstructions as a function of the growth temperature (from [Woolf, 1993a]).

In contrast to the (111)B surface, its (111)A counterpart shows only the (2x2) reconstruction [Cho, 1970;Woolf, 1993b;Woolf, 1993a] (see Figure 3.25).

At high substrate temperature and low As pressure, the GaAs(111)B surface mainly shows the ( √ 19x √ 19) and the (1x1) reconstructions in specific conditions. At lower substrate temperature and higher As flux, it shows instead a (2x2) reconstruction. In contrast, the GaAs(111)A surface displays only the (2x2) reconstruction.

Morphology of GaAs(111)B surfaces

The GaAs(111)B is often found to be non planar. From 550-560 °C, the surface is rough due to pyramids increasing in lateral size as the temperature increases. [Tsutsui, 1990;Woolf, 1993a] (Figure 3.26 (a)). Above 600 °C, these pyramids become more widely spaced [Woolf, 1993a] and good electrical properties are observed [Tsutsui, 1990]. Mirror-like surfaces have been observed at substrate temperatures of 620 °C [Woolf, 1992] and 640 °C [Yang, 1992].

At low temperature, below 550 °C, the surface shows a rather smooth morphology with small pyramids. However, the electron mobility rapidly decreases [Tsutsui, 1990] and no PL signal is observed [Woolf, 1992].

It is observed also that the GaAs(111)B surface roughness increases with the As flux [Yeo, 2000] (Figure 3.26 (b)). With an As 4 flux of 4.9•10 -6 Torr, surfaces show a smooth almost featureless surface [Yeo, 2000]. The morphology changes to round islands with As 4 of 6.2•10 -6 Torr and pyramids at high As 4 (1.1•10 -5 Torr). The electrical properties (electron mobility) also deteriorate with increasing As 4 pressure [Yeo, 2000]. ( from [Yeo, 2000]).

High substrate temperature (620-640 °C) and low V:III ratio are needed in order to obtain a smooth surface morphology and good electrical transport on GaAs(111)B surfaces.

Morphology of GaAs(111)A surfaces

GaAs(111)A surfaces also tend to show pyramidal morphology, but the optimum growth conditions to obtain a smooth surface are different.

Figure 3.27 shows the impact of the substrate temperature and III:V ratio on the surface morphology. Figure 3.27 (a) shows a decrease of the surface roughness with increasing As 4 flux. At lower As 4 :Ga flux ratio, a different type of morphology appears with pyramids rotated of 180°with respect to one another.

A decrease in surface roughness can also be achieved by decreasing the substrate temperature from 650 °C to 550 °C, and it is associated to pyramids decreasing in size (Figure 3.27 (b)). Sato et al. [Sato, 1994] also observed a similar reduction of defect density and defect size, by reducing the substrate temperature and increasing the As 4 flux (Figure 3.27).

Beside these optimum conditions, flat surfaces are hardly obtained. This is attributed to an Ehrlich-Schwobel barrier hindering the incorporation of the Ga atoms at the step edge. These atoms tend to be reflected by the step and to form nuclei on the terraces [Ritzmann, 2018;Esposito, 2017]. A high temperature is necessary to overcome the barrier but the low sticking coefficient of As at high temperature prevents the efficient incorporation. A possible solution is to perform a series of [Ritzmann, 2018]. Summary of the morphology and defect from and resuming graph from Sato et al. [Sato, 1994].

post-growth annealings at 650 °C under an As flux [Ritzmann, 2018]. An alternative route is to use a very high V/III ratio, a substrate temperature of 520 °C [Esposito, 2017] and very low growth rate.

To obtain smooth surfaces and low defect density on GaAs(111)A surface, low substrate temperature and high As flux are required. However, flat layers are not obtained without special procedures, such as periodic annealing between growth steps.

Experiments

In the following, we investigate the surface reconstruction and the morphology of GaAs(111)B layers.

In all the experiments, we set the Ga flux and we vary the As flux and the substrate temperature to find suitable growth conditions. The samples are characterized by Normaski optical microscopy and by AFM to evaluate the surface roughness. Photoluminescence is also used to characterize the optical quality of the layers.

The GaAs (111)B substrates are first degassed at 350-400 °C for 1-2 h is a separate chamber. Once in the growth chamber, the substrates are heated up to 630-640 °C under an As flux to fully deoxidize the surface. The Ga cell temperature is kept at 935 °C, which provide a BEP of 1.44•10 -7 Torr.

Reconstructions observed by RHEED

We first investigate by in situ RHEED the two surface reconstructions of the GaAs(111)B layers.

After deoxidation at T=635-640 °C, the RHEED only shows a 1×1 pattern (Figure 3.28 (a)). After growth of a ∼150 nm thick buffer layer, RHEED shows a √ 19× √ 19 reconstruction (Figure 3.28 (c)) at T=600-640 °C and a (2×2) reconstruction at T=500-520 °C (Figure3.28 (b)). .

We did not observe the intermediate (1×1) reconstruction mentioned in the literature (see section 3.3.1.1), possibly because it only appears in a narrow range of temperature.

Surface morphology as a function of temperature and As flux

We investigated different substrate temperatures (580 to 640 °C) and As:Ga flux ratios (As:Ga BEP ratio from 11:1 to 90:1) in order to obtain smooth layers and good morphology.

The roughness is first evaluated by Nomarski characterization. Some of the samples are then characterized by AFM. The combination of Nomarski and AFM probe the surface at different scale and results in complementary information. For a first indication of the crystal quality, we use photoluminescence (PL).

The grown structure consist in a 90 nm thick GaAs buffer layer, a 12 nm thick Al x Ga 1-x As (x∼0.3-0.4, Al BEP 4-5•10 -8 Torr) barrier layer, a ∼450 nm thick GaAs layer and a 4-6 nm thick Al x Ga 1-x As cap layer. After the growth, the samples are cooled under an As flux.

As flux series

We first characterize samples grown under different As 4 :Ge flux ratios, namely 11:1, 24:1 and 51:1, at fixed substrate temperature of ∼620 °C.

• The AFM images of sample #1 (Figure 3.29), grown under the lowest As flux, shows pyramidal holes of less than 1 µm in width and up to 100-120 nm depth. Similarly, Nomarski microscopy shows small dots related to the pyramidal holes in the AFM image. • At an even higher As flux (As:Ga BEP ratio 51:1, sample #3 in Figure 3.29), the surface morphology is about the same as that of sample #2, with denser atomic steps. From Nomarski microscopy, the surface of sample #3 shows a higher density of conical features compared to sample #2.

Substrate temperature series

We also performed a series of growth experiments at temperatures from 580 °C to 640 °C, with a fixed As 4 :Ga BEP ratio of 42-49:1. The respective AFM and Nomarski images are shown Figure 3.30. The main observations are the following.

• At 580 °C (sample #4), pyramids are visible by AFM and Nomarski microscopy. They are up to 70 nm in height and 2-5 µm wide.

• At 600 °C (sample #5), the pyramids are larger, up to 30 nm in height and 10-15 µm .

• Sample #3 grown at 620 °C, was already shown in Figure 3.29. On large scale AFM images, we observe rather flat pyramidal or conical structures 20-30 µm wide and ∼10 nm high.

• At 630 °C (sample #6), the surface becomes flatter with less conical features in Nomarski images. AFM shows terraces about ∼0.25 µm wide.

• At 640 °C (sample #7), the surface is almost flat. The conical structures are very few and the surface present widely spaced atomic steps, up to ∼0.4 µm wide.

Substrate temperature and As flux

In Figure 3.31, we combine Nomarski characterization of the As 4 flux series with the substrate temperature series and an additional series. The later is a variant of the As flux series but grown at substrate temperature of 600 °C with As 4 :Ga BEP ratio up to 90:1. We also add another sample, grown at high temperature (640 °C) and under high As 4 :Ga BEP ratio (90:1).

For the As 4 flux series at 600 °C, we observe that:

• At low As 4 flux (sample #8) pyramidal holes appears (as for sample #1 of the series at 620 °C).

• For an As 4 :Ga BEP ratio of 21-28:1 or 42-49:1 (sample #9 and #5), pyramidal holes are replaced by pyramidal hillocks.

• For a further increase of the As 4 :Ga BEP ratio to 63:1 and 90:1 (samples #10 and #11), the surface morphology change to dense and well defined small pyramid.

• If we increases the substrate temperature from 600 °C to 640 °C under a high As 4 flux (sample #12), we recover a smooth surface comparable to that of samples #6 or #3 (grown under lower As fluxes at temperature 10-20 °C lower).

Discussion

The pyramids get wider and their facets get shallower with the growth temperature, as observed previously by [Yang, 1992;Tsutsui, 1990]. At smaller scale the atomic terraces become wider with increasing growth temperature and at larger scale the conical features become smoother and wider apart.

At very low As flux we observe pyramidal holes, which disappears if the As flux is increased. With further increase of the As flux, pyramids are formed on the surface which is coherent with observation from Yeo et al. [Yeo, 2000].

Overall, increasing the growth temperature has a similar effect than decreasing the As flux. This is coherent with a higher desorption rate of As from the surface at high temperature. From the experimental point of view, the substrate temperature should be at least 620-630 °C and As:Ga ratio should not be too high (As 4 :Ga BEP ratio < 63:1). However, if the As 4 :Ga flux is decreased too much (BEP ratio < 21:1) we see the appearance of pyramidal holes.

A high temperature substrate (630-640 °C) and a low As flux (21:1 to 49:1) result in the best surface morphology. A too low substrate temperature or a too high As flux result in well defined pyramidal structure. On the other hand, an insufficient As flux induces the formation of pyramidal holes.

Photoluminescence

Description of the experiment Beside the surface morphology, we evaluate the crystal quality using room temperature photoluminescence (PL) using a green YAG laser (532 nm). We characterize samples #4, #5, #3 and #7 from the temperature series (section 3.3.2.2 Figure 3.30). Results On all the spectra (Figure 3.32) we observe a parasitic signal due to the laser harmonic at ∼2λ=1064 nm. The GaAs peak at ∼870 nm is not visible for sample #4 grown at 580 °C. It starts appear at 600 °C (sample #5) and becomes much more intense at 620 °C and 640 °C, which indicates a better crystal quality.

A high substrate temperature (620-640 °C) is needed to obtain the most intense photoluminescence signal.

As dimers versus As tetramers

Description of the experiment GaAs can be grown from two types of As molecules, As 4 and As 2 . The first one is generated with the As cell cracker at 600 °C and the second one with the cracker at 900 °C. We previously studied samples grown using As 4 molecules and we investigate here the effect of As 2 molecules. Growth is performed with the same valve aperture, As 2 :Ga BEP ratio of 35:1 and As 4 :Ge BEP ratio of 46:1. Sample #5 and #8 were previously described in section 3.3.2.2. Samples #13 and #14, grown under As 2 , show no significant difference in surface morphology.

The use of As 2 instead of As 4 does not impact the surface morphology in the studied experimental range.

Conclusion

We have investigated the growth conditions required to obtain smooth GaAs layers on GaAs(111)B substrates. We observe that surface morphology depends on the As flux and on the substrate temperature. From low temperature (580 °C) to higher temperature (640 °C) at constant As flux we observe a collection of pyramidal features which increase in lateral size and decrease in height as the substrate temperature rises. At higher temperature the atomic terrace are clearly visible and widely spaced.

The photoluminescence intensity also increases with the substrate temperature. The GaAs peak appears at room temperature if the substrate temperature is hotter than 600 °C.

At constant growth temperature, increasing the As flux first leads to a flattening the surface by removing the pyramidal holes, before it roughens again due to formation of pyramidal hillocks. before it roughens again due to pyramids hillocks.

The best surface morphologies are obtained at relatively high substrate temperature for growth at 630-640 °C with an As 4 :Ga BEP ratio of at least 21-28:1. No significant change in surface morphology is observed when As 2 is used instead of As 4 .

Heteroepitaxy of GaAs on Ge (111)

In this section we investigate the epitaxy of GaAs layers on monocrystalline Ge(111) substrates. The optimum growth conditions on this substrate will serve as reference for the GaAs epitaxy on the highly (111)-textured MIC-Ge pseudo-substrates.

We first review the literature, starting from the growth on (001) surfaces, then we investigate experimentally the growth morphology as a function of the As flux and the substrate temperature.

The layer are characterized using Nomarski microscopy, AFM, X-ray Diffraction (XRD) and SEM. Specific twins defects are characterized with X-ray pole figures, TEM and EBSD. Finally, we address the question of the polarity of our GaAs layers grown on Ge(111) substrates.

Defects in GaAs grown on Ge(111) -A review

We first shortly review the growth conditions investigated in the literature as well as the main defects encountered.

Growth of GaAs on Ge(100) -Anti-phase domains

When growing a polar material on non-polar substrate, the formation of antiphase domains (APD) has to be considered. We first look at the most common case, that of GaAs grown on Ge (100).

Ge

GaAs

APD boundaries

[100] of GaAs shows incompatible termination of the surface (Ga or As) and they are separated in the layer by specific boundaries where bonds between atoms of the same type are found. In order to avoid the formation of these defective boundaries it is necessary to form only double steps on the Ge (100) surface. This can be obtained by using strong misorientation towards the [011] direction [Pukite, 1987].

According to the literature [Sieg, 1998b;Sieg, 1998a;Ringel, 1997], several measures can be taken in order to avoid APDs in GaAs grown on Ge:

1. add a Ge buffer layer (100 nm thick, annealed at 640 °C) in order to reduce the initial Ge roughness and bury the carbon contamination.

2. using an offcut of 6°toward the [110] in order to form double steps on the Ge (100) surface;

3. avoiding excessive As 2 exposure during the GaAs nucleation 4. cover completely the Ge surface with an As or Ga monolayer.

5. grow the first layer 3 nm of GaAs by migration enhanced epitaxy (MEE) at 350 °C.

Growth of GaAs on Ge(111): full surface steps and termination

Ga As To the contrary of the Ge(100) substrates, the Ge(111) surface should allow the growth of GaAs without anti-phase domains, without the need to use vicinal substrates. The fact that APDs do not appear is due to the fact the surface steps are high of two atomic layers, which implies that neighboring terraces always present the same atomic termination [Morizane, 1977;Patel, 1987] (Figure 3.35).

When exposed to As, Bringand et al. report that the Ge (111) surface changes into a fully As-passivated surface with no dangling bonds [Bringans, 1985;Bringans, [Bringans, 1992] 1987; Bringans, 1992], (Figure 3.36 (a)). This passivated surface results in doubly occupied surface orbitals.

Note that the dangling bonds are eliminated due to the replacement of the outer Ge layer by As atoms. Figure 3.36 (b) shows the density of states for a bulk termination, i.e. a non-reconstructed Ge (111) surface in which the broken bonds correspond to partially unoccupied surface states. With a full As-passivation (Figure 3.36 (c)), the lone pair state is removed and the extra electron of As atoms fully fills the valence band.

Bringans et al. [Bringans, 1987] also shows that there are no stable single monolayer of Ga, starting the GaAs growth on Ge (111) with Ga atoms gives similar results to As passivation. This is explained by the As atoms going below the Ga monolayer to bond with the Ge.

The surface passivations with the highest As coverage are obtained by exposing fully deoxidized Ge(111) surfaces to As 2 or As 4 at 500-600 °C. The obtained Aspassivated surfaces are stable up to 700 °C in vacuum [Bringans, 1992].

Growth of B-polar GaAs on Ge(111)

The As-passivated surface should determine uniquely the polarity of the grown GaAs. Indeed, Koh et al. [Koh, 1999;Koh, 2001] show that only GaAs (111)B can be grown on As-passivated Ge (111). It is even possible to reverse the polarity of a GaAs (111)A layer by inserted a thin Ge layer (3 nm) before resuming the GaAs growth, Figure 3.37, while the opposite is not possible.

Recent works of GaAs on Ge(111) epitaxy also expect, but do not confirm, a (111)B polarity, due to the As surface termination of the Ge(111) [Kajikawa, 2017;Kajikawa, 2018]. [Koh, 2001] Both (111)A and (111)B polarities of GaAs layer on Ge(111) were reported by Bobb et al. [Bobb, 1966] using CVD growth with AsCl 3 and metallic Ga sources in a stream of hydrogen. The GaAs(111)B was obtained in clean conditions whereas Ga(111)A was obtained when the gas stream was contaminated with oxygen. No mixture of the two phases was reported.

A similar situation to GaAs on Ge ( 111) is the case of GaP grown on Si(111). Here, P does not passivate the Si surface as As does [Bringans, 1985]. But Paszuk et al. [Paszuk, 2015] were able to control the polarity of the GaP grown on Si(111) by MOVPE by controlling the passivation of the Si. The standard hydrogen passivation leads to (111)A GaP layers, while the As-passivated Si(111) leads to the GaP (111)B polarity [Paszuk, 2015]. This is a another example of the specific As-passivation of Si(111), initially described by Bringans et al. [Bringans, 1992].

Twins defects in III-V (111) epilayers

While APBs are not likely to occur during growth in the (111) direction, twins can be created. The later tend to form at the Ge/GaAs interface [Kajikawa, 2017;Suzuki, 2013] but they can appear also in GaAs homoepitaxy [Rajkumar, 1991]. More generally, they are common in all zincblende (111) hetero-epitaxy [Kajikawa, 2017]. Twins are created by a stacking fault like defect [Koppka, 2016] (Figure 3.38), which create an armchair-like bonds configuration, instead of the classical zig-zag configuration found in cubic crystals.

A dependence of the twin density with the substrate temperature has been observed in the early days of GaAs epitaxy on Ge(111) using gas sources [Holloway, 1965]. Twinning occurred along multiple planes, especially at low temperature. This early study showed that:

• When grown at 400 °C, GaAs presents a fibre texture.

Twin boundary [1-10] [111] [11-2] [1-10]

[111]

[ stacking-fault like defect • Above 625 °C, twinning only occurs along the <111> axis normal to the surface.

• Between 625 °C and 725 °C (the Ge/As eutectic temperature), some untwinned GaAs layers were obtained, but the results were not reproducible.

Recently, a 100 nm GaSb buffer layer was introduced in order to decrease the twin density for GaAs grown on Ge (111) [Kajikawa, 2017]. The effect is explained by the lower ionicity of GaSb, resulting in a higher energy for stacking fault formation. The volume of twinned domain ratio is reduced from 50% in the GaAs without the GaSb to 20% inside the GaSb and 0.7% in the subsequent GaAs. The twin density is also reduced by the use of As 2 instead of As 4 and by employing vicinal substrates (2°toward [110]).

Twins were also observed in the growth of GaP on Si(111) by MOVPE [Koppka, 2016]. There, the Si surface was terminated with As (see section 3.4.1.2) in order to obtain GaP with (111)B polarity. Their density could be reduced by decreasing the initial growth temperature to 420 °C and by using a 3°off miscut substrate. The twinned fraction can thus be lowered to 8% of the total volume.

Twins have also been observed in the growth of lattice-mismatched GaAs on Si (111) [Suzuki, 2016;Suzuki, 2013]. Suzuki et al. [Suzuki, 2013] show that the twin density can be greatly reduced with an In pre-evaporation, whereas it increases with a Ga pre-evaporation. This group also showed that twins nucleation at the Si interface can be reduced with a lower V:III flux ratio [Suzuki, 2016]. Despite these recent achievements, there is still no clear understanding of how the growth conditions affect the formation of twins formations in GaAs layer grown on Ge(111) surfaces.

We now move to the experimental investigation of the growth of GaAs on Ge(111) substrates.

Growth experiments

MBE growth was performed in a Riber32 equipment. The Ge (111) N-type As-doped substrates (from AXT) were loaded into the growth chamber after being degassed at 350-400 °C in a separate chamber.

The samples were exposed to As 4 for at least 5 min before growth, at 610°C. The GaAs growth was performed at the same temperature using a Ga BEP of 1.4•10 -7 Torr (growth rate of ∼1.5 A/s). For some samples, a thin (Al,Ga)As marker layer was grown with an Al BEP of 5•10 -8 Torr. Arsenic was provided as As 4 , (with the cracker of the cell at 600 °C). Different As:Ga BEP ratios were used, namely 31:1, 62:1, 70:1 and 94:1.

In situ RHEED characterization

Description of the experiment We investigate here the RHEED diffraction patterns of GaAs/Ge(111) grown at different temperatures to compare them with those observed during growth of GaAs layers on GaAs(111)B substrates (section 3.3.2.1 p. 107).

To this end, the Ge (111) surface is first heated at the growth temperature of 600-610 °C and a thin layer of GaAs(10 nm) is grown.

Results Before growth and after the Ge de-oxidation (below 450 °C) the RHEED shows a stable Ge 1×1 pattern up to the growth temperature (Figure 3.39 (a)). After the start of GaAs growth (600-610 °C), the pattern becomes 2x2, Figure 3.39 (c). After the growth of the 10 nm layer, we lower the temperature to 550 °C and the pattern remains 2x2, Figure 3.39 (b). Heating the surface up to 635 °C does not change the pattern (not shown). The pattern is always streaky, which is indicative of a smooth surface. Discussion The (1x1) RHEED reconstruction of the Ge(111) surface after deoxidation is characteristic of the Ge surface [Koh, 2001;Kawai, 1993]. For GaAs (111)B layers, we should observe a 2×2 pattern at low temperature (less than 550 °C) [Woolf, 1993a] and a √ 19x √ 19 reconstruction at high temperature (560-620 °C). In this temperature range we observed instead a clear and stable 2×2 pattern, which is compatible with a (111)A orientation.

The growth of GaAs on Ge produces a planar surface. We observe the 2×2 surface reconstruction at high and low temperature, in contrast to the expected √ 19x √ 19 structure at high temperature for GaAs(111)B layers. The RHEED pattern is compatible with the GaAs(111)A orientation.

Ex situ AFM

Description of the experiment To confirm the flat morphology observed by RHEED, we perform AFM on the thin GaAs layer (∼10 nm). We compare this thin layer to a thicker one (∼450 nm) grown in the same conditions. Figure 3.41 shows the reciprocal space maps around the 111 and 331 reflections. In both maps, the horizontal corresponds to an in plane direction and the vertical one to the substrate normal. The alignment of the GaAs and Ge peaks in the vertical direction proves that the GaAs layer is fully lattice-matched to Ge in the (111) growth plane (i.e. "fully strained").

We also observe Pendellösung fringes which match the layer thickness estimated from the previous RHEED calibration.

HRXRD shows that the GaAs epitaxial layer is fully strained to the Ge (111) substrate.

As flux and substrate temperature

Description of the experiment To decrease the surface roughness observed in the experiments described in the previous section, we now vary the As flux (at constant Ga flux) and the substrate temperature.

We grow here ∼450 nm thick GaAs layers terminated by a thin Al x Ga 1-x As (6 nm) layer protected by a thin GaAs (4 nm) cap layer (x∼0. 3-0.4, Al BEP 4-5•10 -8 Torr). The Ga BEP is the same as in the previous section (∼1.45•10 -7 ), yielding a growth rate of GaAs of ∼ 1.5Å/s. As 4 is varied to get V:III BEP ratios of 37:1, 63:1 and 94:1. We consider two growth temperatures, 600-610 °C and 630-640 °C.

Samples are characterized ex situ by Nomarski microscopy and SEM. We now characterize the same samples with SEM in top view and tilted (45 °) view (Figure 3.43).

Result

• sample A: pyramidal roughness with two possible orientations of the pyramids, rotated 180 °apart, see (Figure 3.45)

• sample B: relative smooth surface small pyramids (0.2-0.5 µm)

• sample C: large terraces separated by a rough ridge.

• sample D: pyramidal roughness; single orientation of the pyramids in the selected area

• sample E: relatively smooth surface with small pyramids • sample F: the surface structure appears complex • At low As 4 flux, the surface shows a pyramidal roughness, for growth in the range 600-640 °C.

• At intermediate As 4 flux, the surface roughness is much lower.

• At higher As 4 flux, the sample is rough but without pyramids. The surface is smoother at higher temperature, but its structure is complex.

Pyramids and hillock Nucleation study: Ga or As first

We have seen that at the early stages of growth, the GaAs layer is flat, whereas it displays large pyramidal features for longer growth duration. To find out when these pyramids form, we grow here a ∼45 nm thick layer. We also investigate if the nucleation of GaAs on Ge is affected by the As or Ga pre-exposure (up to 5 min for As Results Figure 3.46 shows SEM cross-sectional view of a ∼45 nm thick GaAs layer grown on Ge using As 4 pre-exposure. We observe pyramids with high aspect ratio, about twice thick than the surrounding GaAs layer. The pyramid and their facets are inclined by about 22°to the horizontal (in this particular SEM projection). Discussion The As passivation of the Ge surface before growth should be a key parameter to obtain untwinned GaAs(111)B [Bringans, 1985]. Thus, we are not surprised that starting the growth with Ga result in defective structures. However, this point would require further investigation.

Using As 2 instead of As 4 is expected to yield a lower twin density [Kajikawa, 2017]. Here, we only observe a minor effect on the shape and density of the pyramids.

Pyramids appears already for thin GaAs layers (45 nm) Starting growth after Ga pre-exposure induces morphological defects. Using As 2 instead of As 4 has only a minor effect on on the layer morphology.

Flattening the GaAs layer by MEE

We investigated the use of Migration Enhanced Epitaxy (MEE) [Takano, 1991] to reduce the roughness of the ∼45 nm thick GaAs layers on Ge(111). We use two type of MEE sequence:

• In our "type-I" MEE experiments, we expose repeatedly the layer to As only for 4 s and to Ga only for 4 s, for a total growth duration of 5 min.

• The MEE "type-II" sequence consists in repeated exposure to single element Ga and As 4 for 2 and 4 s, respectively, followed by 2 s without any flux. We grow 76 cycles. 3.48 (c) shows the sample with the type-II MEE layer. We observe the appearance of additional small pyramids but the morphology is similar. Hence, only a marginal flattening effect could be obtained by using MEE. However, our system is not well adapted for such type of growth, since a minimum time of 2 s between the opening and closing of a shutter is required in order to preserve the delicate mechanics of the machine. Due to these practical difficulties, we do not further investigate the MEE methods.

Results

Within the limitations of our MBE setup, we find that MEE has little effect in flattening the layer.

Impact of layer thickness on surface morphology

To get a broader view of how the roughness evolves with epilayer thickness, we now study an even thicker layer.

Results Figure 3.49 shows SEM top view, cross-sectional view and 45°-tilted view of GaAs grown on Ge(111) with thicknesses of ∼45 nm, ∼450 nm and ∼1.63 µm. The As 4 BEP is ∼7•10 -6 Torr for the first two samples and ∼1.1•10 -5 Torr for the third one. Substrate temperature is 600-620 °C. The layers become flatter with increasing thickness, despite the initial formation of pyramids. Independently of the formation of pyramids at the initial stage of the growth, it is possible to obtain smooth layer at 600-620°C by increasing the growth duration and the epilayer thickness.

Defect characterization

In several cases, we have observed two type of pyramids, the second being rotated by 180°from the first (see Figure 3.45). We investigate here this type of defects using XRD, TEM and EBSD. We observe evidence of twinning in the epitaxial GaAs layer.

The twinned domains are very important at low and high As 4 fluxes, with respectively single and multiple order twinning.

Only the sample growth using intermediate As flux (BEP 8-9•10 -6 Torr) shows little sign of twinnig.

Discussion

In the TEM cross-section, the twinned crystal appears to start in the middle of the layer. However it is not clear if it is really the case or if the twin domains initially started at the GaAs/Ge interface and later expanded in three dimensions during the growth. In the later case, the current FIB cut may not have sampled the initial twin location.

TEM cross-section view confirms that the area with rotated pyramid correspond to twins inside the material thickness. No other defects are visible in the TEM cross-section in the twinfree areas.

EBSD analysis

Description of the experiment In order to unambiguously characterize the spatial extension of the twinned fraction of the sample, which could be guessed from SEM contrast in Figure 3.51, we perform EBSD on the sample grown using the optimized sample made with intermediate As flux (sample B (77101) of section 3.4.2.4, ∼450 nm thick GaAs with (Al,Ga)As/GaAs cap layer) . Results Figure 3.54 shows SEM and corresponding EBSD maps. We check that the sample is (111)-oriented. Some small zones show a slight misorientation with respect to the EBSD detector, indicated by arrows in Figure 3.54 (a,b). Further analysis reveals that they correspond to twinned domains, shown by Euler map (Figure 3.54 (e)) and the zoomed EBSD map with the projected unit cell orientations (Figure 3.54 (d)).

Discussion

The EBSD analysis confirms that only a small fraction of the sample surface is twinned. Considering the previous TEM analysis, this indicates that the majority of the volume of the sample is twin-free. It also ascribes the origin of the SEM contrast (Figure 3.51) to the different crystal orientation of the twinned domain.

EBSD confirms that the sample is mostly twin-free and that only the small areas with rotated pyramids are associated with twinned domains.

Cathodoluminescence

Description of the experiment We characterize a sample similar to the one used for EBSD, but with no (Al,Ga)As/GaAs cap layer. We use here cathodoluminescence (CL) to better understand the sample structure. Compared to EBSD, this technique can inject electron-hole pairs in the full thickness of the layer (450 nm) to determine the spatial extension of twin boundary from top view images. In Figure 3.55 (c,d), we focus on a defective zone with a complex twin boundaries. In Figure 3.55 (c), the red (α) and blue (β) areas mark the zones related to the two possible in-plane orientation, using the CL image as a guide. Figure 3.55 (e) shows the cross-section schematic of the interface between the two twinned phases. The purple dotted line shows the incoherent interface across the layer thickness which holds non-radiative recombination centers. The blue dotted line instead shows the coherent interface in the plane between the phases. We observe that the center of the pyramidal defects are highly luminescent even if they are part of a twinned area, possibly due to the increased thickness.

Discussion

The vertical boundary between two twinned areas hold a large density of non-radiative recombination centers. The corresponding CL signal drops markedly in the vicinity of this the interface. This poor crystal quality is explained by the incoherent interface formed in the vertical direction after a twinning across the growing (111) plane, as observed in the TEM characterization, (see section 3.4.3.2, p. 133).

We observe that the dark CL zone does not extend much further than the top surface associated to the rotated pyramids. This also confirms the previous TEM observation, in which the twinned crystal location was limited to zones close to the top area identified with rotated pyramids. This supports that the vast majority of the GaAs volume is twin-free.

The interface between two twinned domains is characterized by a significant decrease of the CL intensity, consistent with an incoherent twin boundary across the GaAs thickness. The relatively small extension of the low intensity CL zones indicates that the twinned domains are limited to the area marked with rotated pyramids and that the vast majority of the GaAs layer is twin free.

Epilayer polarity

Description of the experiment In all the previous characterizations, we never observed antiphase domains. It is very likely that GaAs presents a single polarity. In the literature, Koh et al. [Koh, 2001] have shown that GaAs grown on Ge(111) is expected to show the (111)B polarity due to the As-termination of the Ge(111) surface [Bringans, 1985].

In this section, we perform HRSTEM HAADF imaging, coupled to atomic resolution EDX spectroscopy to determine the polarity of our GaAs layers.

Results Figure 3.56 (a) and (b) shows two schematics of the crystal structure of GaAs(111)B and GaAs(111)A layer respectively. We note that GaAs(111)B is terminated by As atoms, while GaAs (111)A is terminated by Ga atoms. Figure 3.56 (c) shows a high-resolution high-angle annular dark-field imaging (HAADF) of the growth GaAs layer. The HAADF image intensity is typically proportional to Z 2 , where Z, the atomic number of the imaged atomic column. We find Z=31 and Z=33 for Ga and As, respectively. Thus, there should be a small contrast between Ga and As atomic columns. away from the substrate, which suggests that the topmost atoms is Ga. In this configuration, the GaAs layer orientation is the GaAs (111)A polarity.

To confirm this polarity, we performed atomically resolved STEM-EDX analysis on a small area (Figure 3.57 (a)). The corresponding HAADF image is shown in Figure 3.57 (c) and the schematics of the crystal structure in Figure 3.57 (b). The As atoms of each dumbbell are clearly at the bottom side, i.e the substrate side, which confirms the GaAs layer grows in the (111)A polarity. Discussion Our results do not match the work of Koh et al. [Koh, 2001] who found a (111)B polarity for GaAs grown on Ge(111). This finding also challenges the fully As-passivated ideal Ge (111) surface advertised by Bringans et al. [Bringans, 1985] for our samples and points toward possible contamination of the initial Ge (111) surface by oxygen or other compounds [Bobb, 1966].

HAADF As Ga 5 Å 5 Å [1-10] [11-2] [-1-1-1] (a) (c) (b) (d)
Retrospectively, the observed GaAs (111)A on Ge (111) polarity allows us to better understand the different surface morphology between the GaAs(111)B reference layers (see section section 3.3.2.2 p. 108) and the GaAs layers grown on Ge (111) substrates (see section 3.4.2.4 p. 124). For GaAs (111)A layers, twinning can occur not only at the interface between the Ge (111) substrate and the GaAs (111)A layer, but at any time during growth [Rajkumar, 1991;[START_REF] Park | Analysis of twin defects in GaAs(111)B molecular beam epitaxy growth[END_REF], depending on the As flux.

For high As fluxes, we observe multiple order twinning, while the literature on GaAs (111)A homoepitaxy reports a smoothing of the surface and no visible twin defect [Sato, 1994;Esposito, 2017]. It is therefore likely that the high As:Ga ratio induces twin defects at GaAs/Ge interface, which later propagate in the layer. Conversely, the twins observed at low As flux may be generated at any time inside the layer, not only at the initial GaAs/Ge interface.

The GaAs layers grown on Ge(111) show a (111)A polarity, contrary to what is reported in the literature. This allows one to reinterpret the dependence on the surface morphology with the As flux. At high As flux, twins are most likely generated at the GaAs/Ge interface and later propagate in the layer. At low As flux twin can generate anywhere in the layer. At intermediate flux, we find a narrow regime for which the twin are not generated at the starting GaAs/Ge interface and in the GaAs layer.

Interdiffusion

Description of the experiment It is known that the atoms from GaAs and Ge can interdiffuse [Kawai, 1993;Kawai, 1992]. To estimate the interdiffusion length, we performed large scale EDX spectroscopy on cross sectional TEM specimens.

Results Figure 3.58 (a) shows an HAADF image extending from the Ge wafer to the GaAs layer. We clearly observe the thin (Al,Ga)As (8-9 nm) layer above the 90 nm GaAs buffer layer. This experiment shows that there is some diffusion of Ge into the GaAs, over 20-30 nm from the interface. Similarly, we find both Ga and As into Ge but over a reduced length (5-10 nm).

We report also the EDX map for each element (Ge, As, Ga, Al) in Figure 3.58 (d-g) showing up the spatial distribution of the concentration of each material in each layers.

Discussion Ge diffuses more into the GaAs layer over a distance larger than that over which Ga and As diffuse into the Ge substrate. This agrees with previous reports from the literature [Kawai, 1992;Kawai, 1993]. This Ge diffusion may result in the n-type doping of the GaAs layer, which is not problematic for our future solar cells design. On the Ge side, we only use ntype Ge substrates for which the co-diffusion of As and Ga may lead to possible compensation. Ga and As diffuse into the Ge substrate over about 5 to 10 nm and Ge diffuse into the GaAs layer over about 20 to 30 nm.

Conclusion

In this section, we define the optimum conditions for the growth of GaAs on Ge(111).

We have shown that the growth of GaAs on Ge in the (111) direction is affected by twinning. The twinned domains are characterized by a "vertical" non-coherent interface and an "in-plane" stacking fault, which we observe by TEM. The defective interface is responsible for the drop of luminescence in the GaAs layer measured by CL.

We also showed that by adjusting the growing conditions, the twin density can be reduced during the direct growth of GaAs on Ge(111), without any buffer layer.

The adjustment of the As:Ga flux ratio plays a critical role in avoiding twin formation. Only a narrow range of fluxes has been found suitable (for the given growth temperature of 610 °C)

Contrary to what could be expected from the literature, we demonstrated that GaAs layers grown on Ge( 111) have (111)A polarity. This allows one to relate the twin formation at low As 4 flux to unsuitable growth conditions for GaAs(111)A layer and the twin formation at high As 4 flux to twinning of the initial GaAs nuclei on the Ge(111) surface. These optimized growth conditions will serve as reference for the future growth of GaAs layers on MIC-Ge pseudo-substrates.

Growth of GaAs on the MIC-Ge pseudosubstrates

In this section we work towards the epitaxial growth of GaAs on MIC substrates, with a particular focus on the selective growth of GaAs on the crystallized Ge patches with respect to the silica surface.

We investigate how the GaAs surface roughness evolves with the thickness of the epilayer in comparison to the reference growth experiments on monocrystalline Ge(111) substrates (section 3.4.2.6, p. 130). We finally consider the crystalline quality of the grown GaAs layer and the associated crystal texture with respect to the underlying MIC-Ge template.

Selective growth with respect to silicon dioxide

Surface cleaning

Description of the experiment To achieve a good selectivity, the silica surface should be as clean as possible. After the Al wet etching (H 3 P O 4 at 50-80 °C), we clean the Ge surface by successively dipping the sample in a H 2 O 2 (30 s) and HF (30 s) solutions, rinsing the surface with deionized waters between each treatment. This full cleaning sequence is repeated twice.

H 2 O 2 is expected to etch the Ge residues and the chemisorbed carbon from the surface [Amy, 2006]. HF etches the first nanometers of the silica and also removes the GeO 2 native oxide [Deegan, 1998]. The other sub-oxides are partly removed during the deonised water rinsing [Onsia, 2005]. Note that our 1% HF solution cannot passivate the Ge surface with hydrogen, as this requires concentrations over 10% [Amy, 2006].

To investigate the selectivity of GaAs growth, we use non-patterned MIC-Ge substrates with a low Ge coverage of the silica.

Results Figure 3.59 (a,b) shows 45°-tilted SEM images of MIC-Ge islands on the silica after the wet etching of the Al (with ortho-phosphoric acid). Figure 3.59 (c,d) shows the same sample after the full cleaning sequence using H 2 O 2 and HF. The residue on the silica surface, Figure 3.59 (a,b) are not visible using SEM after the full cleaning procedure.

The simple Al wet etching (using H 3 PO 4 ) leaves residue on the silica surface between the MIC-Ge crystals. An additional cleaning procedure using H 2 O 2 and HF is required to obtain a cleaner silica surface.

Ge

Growth temperature and selectivity

Description of the experiment We investigate different growth temperatures from 600 °C to 670 °C, in order to find the optimum conditions for selective growth. The As 4 BEP is ∼7•10 -6 Torr and the Ga BEP is ∼1.5•10 -7 Torr.

The MIC substrates are cleaned with H 2 O 2 and HF as detailed in the previous section. The samples are degassed at 450-500 °C before being entering the growth chamber. The samples are exposed to As for 5 min at the growth temperature before starting growth. We grow ∼360 nm of GaAs, capped with a thin (Al,Ga)As (10 nm) and GaAs (4 nm) layers.

Results Figure 3.60 shows that at 600 °C and 620 °C the GaAs growth is not selective and proceed similarly on the MIC-Ge and on the silica. At 635 °C, the growth becomes potentially selective but we still observe small Ge nuclei on the silica. At 645 °C, these nuclei persist and but the GaAs growth on the MIC-Ge is not continuous anymore. At 670 °C, we only observe isolated grains of GaAs. Large fractions of the MIC-Ge templates are missing between the GaAs grains.

Discussion

To achieve partial selectivity the growth temperature should be higher than 620 °C. However, parasitic GaAs nucleation on the silica persists even at 645 °C. At higher temperature, the GaAs growth is not continuous anymore and the MIC-Ge pads start to shrink.

Therefore, there is no temperature for which we can obtain perfect selectivity on the silica without damaging the underlying MIC-Ge layer. The best compromise is found near 635°C, with a reasonably flat GaAs layer and only partial selectivity.

GaAs growth selectivity on the silica starts from 620-635 °C. However, parasitic nucleations on the silica persist even at higher temperature, despite the dedicated surface cleaning. Above 645 °C, the MIC-Ge layer starts to decompose during the GaAs growth and the later is not continuous anymore.

Layer thickness

Description of the experiment As we see in the previous section, the surface of the ∼360 nm of GaAs layer grown on MIC-Ge is very rough. We performed a series of growth experiments with the different layer thicknesses to characterize the evolution of this surface roughness. As 4 BEP is 6-7•10 -6 Torr for the sample thicknesses up to 720 nm and 1.1•10 -5 Torr for the thicker sample. Substrate temperature is 625-640 °C.

Results

The SEM top, side and 45°-tilted views of Figure 3.61 show that the sample with the thinnest GaAs layer (∼45 nm) has the smoothest surface. The roughness increases with the layer thickness for all layers. For thickness larger than 720 nm, the selectivity on the silica surface is lost.

From the cross sectional view we observe a peak-to-valley roughness of ∼80 nm for the ∼45 nm thick sample; ∼170 nm for the ∼360 nm thick sample; ∼300 nm for the ∼720 nm thick sample and ∼750 nm for the 2 µm thick sample.

The series in thickness can be grown with As 2 (BEP of ∼4•10 -6 Torr) instead of As 4 , Figure 3.62, with comparable results.

Discussion

As opposed to the case of GaAs growth on Ge(111) single crystal substrates, for which the top surface can be made flatter with increasing layer thickness (see section 3.4.2.6, p. 130), the GaAs layer grown on the MIC-Ge template only becomes rougher with increasing layer thickness.

The GaAs layers grown on the MIC-Ge pseudo substrates cannot be smoothed by increasing the total GaAs thickness, independently of the As source (As 2 or As 4 ).

Ex-situ characterization

Description of the experiment We combine now SEM imaging with secondary electrons and EBSD characterization. The goal is to determine the quality (in terms of crystal orientation and roughness) of the GaAs material grown on the MIC-Ge and how this relates to that of the MIC-Ge pseudo-substrate. To avoid issue related to large surface roughness, we study a 45 nm thick GaAs layer grown on MIC-Ge using As 4 (BEP of ∼9•10 -6 Torr) at a substrate temperature of 630-640 °C.

We characterize the crystalline orientation of the grown GaAs layer using Electron Backscatter Diffraction (EBSD) and we compare it to that of the MIC-Ge pseudosubstrate prior the growth. If the GaAs growth proceeds by epitaxy, we should observe the same crystal textures in both layers.

Results Figure 3.63 shows low magnitude images of the MIC-Ge layer before (upper row) and after (lower row) the GaAs epitaxy. The images of each row correspond respectively to SEM and to the crystalline orientation projected in the vertical direction and in-plane, as determined by EBSD. Detailed views of the same sample are presented in Figure 3.64.

The large area views show that the MIC-Ge (Figure 3.63 (b)) and the epitaxial GaAs (Figure 3.63 (c)) are both highly (111)-textured, with small isolated (100) crystals. Similarly, the in-plane crystal orientation is comparable before and after growth, with large domains of uniform random in-plane orientation arranged around a central dendrites (Figure 3.63 (c,f)). This type of structure is characteristic of the MIC-Ge crystallization in our samples, as shown in section 2.6.3 ( chapter 2 p. 62).

The detailed views of Figure 3.64 confirm the above observations. We observe a clear (111) crystal texture, before the growth (MIC-Ge substrate), Figure 3.64 (b) and on the grown GaAs layer, Figure 3.64 (e). With the exception of the (100) oriented defects, the in-plane orientation of the (111)-textured crystal is the same all along the characterized area.

Discussion

The clear conservation of the (111) crystal texture and the in-plane arrangement is the definite proof that the GaAs growth proceed by epitaxy on the MIC-Ge pseudosubstrate. This is only possible if the Ge surface is oxide free and clean enough to permit the coherence of the two crystal lattices. The GaAs layer are grown epitaxially on the MIC-Ge pseudosubstrate and conserve the initial (111) crystal texture and the in-plane geometry.

Conclusion

In this section we succeeded in growing GaAs epitaxially on a MIC-Ge pseudosubstrate. This GaAs material has the same (111) texture as the MIC-Ge substrate. Despite this achievement, which indicates a rather clean Ge surface before the GaAs growth, several significant challenges remain unanswered.

The GaAs growth selectivity on the MIC-Ge pseudo substrates is only partial at 635°C. After the deposition of ∼500 nm layer, randomly orientated GaAs crystals nucleate on the silica surface and enlarge with further growth. Any lower substrate temperatures reduce the growth selectivity and any higher temperatures damage the MIC-Ge patches during the GaAs growth.

The roughness of the GaAs layer increases with the layer thickness. Above 45 nm thick, the obtained GaAs surface is so rough that it cannot be properly characterized by EBSD anymore. This raises important concern for the growth of full solar cell 3.6. Conclusion on GaAs and Ge epitaxy structures, which are about 1-2 micron thick.

Conclusion on GaAs and Ge epitaxy

The goal of this chapter was the optimization of GaAs epitaxy on MIC-Ge substrates. Because of the high (111)-texture of the MIC-Ge observed in chapter 2, we decided to optimize first the homoepitaxy of GaAs on GaAs(111)B substrates and second the heteroepitaxy of GaAs on single-crystal Ge(111) substrates.

On GaAs(111)B monocrystalline wafers, we showed that a high temperature (630-640 °C) and a low As 4 flux results in a flat surface morphology and a high luminescence intensity at room temperature.

On Ge(111) monocrystalline substrates, twin defects can form in the GaAs layer. Using CL and TEM, we show that the twinned domains are separated by grain boundaries with hold large densities of non-radiative recombination centers. We show that the twin density can be minimized by adjusting the growth conditions. The As:Ga flux ratio plays a critical role in avoiding the formation of twin and only a narrow range of As flux (As 4 BEP close to ∼9•10 -6 Torr) and growth temperature (610-640 °C) is found to produce defect-free layers.

Rather surprisingly, we observe that the GaAs layers grow in the (111)A polarity on Ge(111), whereas the literature mostly report (111)B layers. Still, these GaAs (111)A layers can present a smooth surface if the layer thickness is large enough (∼500 nm). Tentative smoothing of the initial stage of the growth using MEE or variant thereof did not produce any significant improvement.

Concerning the MIC-Ge substrates, we first investigated the growth of a Ge buffer layer. Despite the good deoxidation of the Ge surface (followed by XPS), the Ge growth selectivity is only partial and the Ge layer only forms large grains, which significantly increases the surface roughness.

The direct growth of GaAs on the MIC-Ge pseudo-substrate is also partially selective but compared to Ge, the GaAs growth produces a uniform coverage of the MIC-Ge templates. We confirm the GaAs growth proceed by epitaxy and that the (111) preferential orientation of the MIC-Ge is transferred to the GaAs layer.

However, we also observe that the GaAs layer on MIC-Ge becomes increasingly rough with thicker layers. For thickness of approximately two micron, comparable to that of solar cell structure, the final surface is very rough (> 500 nm) and the initial growth selectivity is lost on the silica surface. This raises significant concerns for the growth of solar cell structures on the MIC-Ge pseudo-substrates.

Introduction

The aim of this final chapter is to fabricate full solar cell structures based on a GaAs absorber, using the previously optimized MIC-Ge pseudo-substrates. Before the epitaxy on this complex substrate, we first perform the MBE growth of GaAs solar cell structures on reference monocrystalline substrates: GaAs(111)B, GaAs(111)A and Ge(111) (Figure 4.1), to test and optimize the device performance.

In this work, we use a relatively simple AlGaAs/GaAs/AlGaAs heterostructure for the active region of the solar cell device. This structure requires no effort in lattice match between the different materials, GaAs/AlGaAs, so that it can be easily transferred from one substrate to another, independently of the possible temperature variations due to different doping, thicknesses or crystal structure of the used substrates. .

Top contact

GaAs N+ wafer substrate

GaAs solar cell

Back contact

Si/SiO 2 \ SiO 2 The solar cells were grown either in the Riber 32 MBE described in chapter 3 (section 3.2.2, p. 98) or in the RIBER Compact 21 MBE, see picture in Figure 4.2. This last machine is connected to the Si-Ge MBE chamber, previously described in chapter 3, (section 3.1.2.1 p. 86). [Wang, 2017] and Venkatasubramania et al. [Venkatasubramanian, 1996]. Both cells have a similar structure, with a buffer/contact layer, a back-surface field (BSF) layer, a thick n-doped GaAs base, a thin p-doped GaAs emitter, a window layer and a final p+ top contact layer.

Epi-layers

Top contact

Ge N wafer substrate

GaAs solar cell

Back contact

Top contact

MIC-Ge pseudosubstrate

Back contact

GaAs solar cell

In Figure 4. 3 (b) we observe that a GaAs undoped spacer layer has been inserted between the n-type base and the p-type emitter. This layer is shown to reduce the dark current generated at the grains boundaries and to limit the reduction of the open circuit voltage (V o c) which occurs when defective growth is performed on polycrystalline Ge substrates [Venkatasubramanian, 1996]. This p-i-n variant could be an interesting candidate for our MIC-substrate, due to the proximity between the poly-crystalline Ge substrates and the MIC-Ge pseudo-substrates.

On (100) silicon, 11.8% efficiency was obtained by Wang et al [Wang, 2017] (without antireflection coating), despite the defects generated in the Si x Ge 1-x graded buffer layer (short circuit current J sc = 19.56mA/cm 2 and open circuit voltage V oc = 0.9 V). On poly-crystalline Ge substrates, Ventkatasubrama et al. [Venkatasubramanian, 1996] obtained a 18.2% efficiency (using antireflection coating), despite the defective grain boundaries in the active GaAs layer (J sc = 23 mA/cm 2 and a V oc = 0.99 V).

Venkatasubramanian et al. also report the beneficial use of Se or S as a n-dopant in the increase of the short circuit current in the presence of grain boundaries [Venkatasubramanian, 1996]. In the current MBE configuration, these dopants are not available for our study and we will only use Be (p-type) and Si (n-type/p-type) impurities.

Doping of GaAs (111)A and (111)B layers

State of the art

The doping of the PN junction is a critical requirement of the solar cell. As the incorporation site and the resulting doping type may depend on the crystal orientation, we focus our work on the doping of GaAs (111) A and (111)B layers, with a short literature review complemented by our own experiments. 111)B (black full circle) and GaAs(100) (empty circles) as a function of the temperature of the Si effusion cell (from [Woolf, 1992]).

GaAs(111)B

A systematic study of the Si doping of GaAs layers grown on GaAs(111)B substrate has been made by Woolf et al. [Woolf, 1992] and compared to that of reference GaAs(100) layers (Figure 4.4). The n-type carrier densities measured in the GaAs(111)B layers is always marginally larger than the corresponding value in GaAs(100) layer. Nevertheless, the behavior of the two crystal orientations are comparable for this particular range of carrier densities (6•10 14 to 10 18 cm -3 ).

The small variation between the two crystal orientations may be related to defects with donor like behavior for the GaAs (111)B layers [Woolf, 1992]. This is coherent with the higher electron mobilities measured in GaAs(100) layer compared to that measured on GaAs(111)B for comparable carrier densities As far as the Be-doping of GaAs (111)B layers is concerned, we could not find any recent reference study on it. Works from Pabla et al. [Pabla, 1993] seems to indicate that p-type doping can be achieved. However, this remains to be confirmed by recent experimental work.

N-type doping of GaAs(111)B layer is possible with silicon impurities, with minimal adjustments from the growth conditions of reference GaAs (100) layers. P-type doping of GaAs(111)B may be possible using Be impurities but experimental verification is required. [Sato, 1996]). Degree of compensation of the Si impurities in GaAs layers, (111)A (circle), ( 111)B (triangle) and (100) (cross) (b), extracted from [Yaremenko, 2008].

GaAs(111)A

In contrast to the GaAs(111)B shown above, GaAs(111)A layers show a different doping type with Si depending on the growth conditions [Sato, 1996]. The later is an amphoteric impurity and can act as a n-type dopant (donor impurities), typically at low substrate temperature and high As 4 :Ga ratio or as a p-type dopant (acceptor impurities) in the opposing conditions, see Figure 4.5 (a).

The final value of doping concentration is the result of the total compensation in the layer. At high As fluxes the Si mainly occupy Ga sites, which results in the n-doping of the layer. Conversely, the Si mostly occupy As sites at low As fluxes, which results in p-doping [Yaremenko, 2008]. Yaremenko et al. , define a degree of compensation, k = N a /N d , which is unity when the layer is fully compensated [Yaremenko, 2008]. At low As flux, they observe that the (111)A layer is close to full compensation, Figure 4.5 (b) circle, while the (111)B and (100) layers are only marginally compensated. The n-type doping behavior is restored at higher As fluxes. This particular behavior of the (111)A layers with respect to (111)B or (100) layers could be related to the relatively low sticking coefficient of As 4 tetramers on this surface [Sato, 1996].

We could not find any reference of the Be-doping of GaAs(111)A layer. Experimental work are required to check the doping and the associated carrier density.

Si impurities show an extreme amphoteric behavior in GaAs (111)A layers. N-type doping is possible at low temperature and large As fluxes, far from the typical grown conditions of GaAs (100) or (111)B layers. The Si-doped GaAs(111)A layers may still suffer from large compensation, even in these specific growth conditions. The Be-doping of GaAs(111)A layer is not documented.

Experiment on GaAs(111)B

We first investigate the doping of GaAs (111)B layers using Si as the n-type dopant and Be as the p-type dopant. We use a test growth structure in which the layer to be characterized is protected by two thin AlGaAs layers from the substrate and the air oxidation.

The test structure consist in: ∼90 nm undoped GaAs buffer layer; ∼12 nm of Al 0.4 GaAs undoped; ∼450 nm of doped GaAs; ∼5 nm of doped Al 0.4 GaAs and ∼4 nm of doped GaAs cap layer. The schematic of the structure is shown in Figure 4.6. The value of carrier concentration and mobilities are measured by Hall effect.

Silicon doping

Description of the experiment We compare samples with variable Si doping, grown using Si cell temperatures of 970 °C, 1040 °C and 1110 °C. The substrate temperature is maintained constant at ∼600 °C. The As flux consists in As 4 tetramers (As cell cracker temperature 600°C).

This first series is completed by two other samples grown at intermediate Si flux, first using a higher substrate temperature (640 °C), and second using high substrate temperature and As 2 dimers (As cell cracker temperature 800 °C).

The As 4 BEP is about (6.5±1)•10 -6 Torr. The As 2 BEP is ∼4.5•10 -6 Torr. The Ga BEP flux is about (1.5±0.05)•10 -6 Torr, equivalent growth rate of ∼1.5 Å/s.

Results

Figure 4.7 shows the n-type doping of the GaAs (111)B layer as a function of the Si cell temperature. The carrier density follows an exponential trend with the inverse of the temperature of the Si cell (Arrhenius law), which means that it is roughly proportional to the impinging Si atomic flux.

We observe that the use of As 2 instead of As 4 or the increase of the substrate temperature to 640 °C have no significant impact on the doping concentration.

As a reference, Figure 4.7 shows also the doping value calibrated for a GaAs (100). 111)B compared to reference GaAs(100). Substrate temperature is ∼600 °C (black squares) and ∼640 °C (green circle and red triangle). Square and circle samples are grown using As 4 , while the triangle sample is grown using As 2 .

We confirm the n-type doping of Si impurities in GaAs (111)B layers, with carrier densities similar to that of GaAs (100) layer grown in comparable condition.

Beryllium doping

We perform the similar experiment with Be. Figure 4.8 shows the p-type carrier densities of GaAs (111)B layers as a function of the Be cell temperature. As in the case of Si, the doping concentration follows an exponential trend with the inverse of the temperature of Be the cell, which indicates that is it proportional to the Be atomic flux. As a reference, Figure 4.8 shows also the doping value calibrated for a GaAs (100). We confirm that the p-type doping of GaAs(111)B layer can easily be made using Be impurities, in growth conditions comparable to that of GaAs (100) reference layers.

Effect of the Si and Be doping on the (111)B surface morphology

Results Figure 4.9 shows Nomarski characterization of a reference GaAs (111)B undoped layer (section 3.3.2.2, p. 108) compared to Si-doped (Si at 1060 °C) and Be doped (Be at 750 °C) layer. We observe that the Si doped sample shows the same morphology of the undoped sample, while the Be doped sample presents a relatively smoother surface.

Discussion

The smoothing behavior observed with Be can also be observed without Be using higher substrate temperature (section 3.3.2.2, p. 108). While we cannot rule out a possible drift of the substrate temperature for the Be doped samples, this may indicate a possible surfactant effect of Be impurities on the GaAs(111)B surface.

Experiment on GaAs(111)A

The experiment on GaAs(111)A doping could not be performed due to limited access time on the MBE machine.

Conclusion

The Si doping of GaAs(111)B is well documented and shows low variation with respect to Si doping of GaAs(100). On the contrary few literature on Be doping could be found. In this section, we experimentally verified that Si and Be of GaAs(111)B are respectively efficient n-and p-type impurities.

For the GaAs(111)A, Si is known to incorporate either in Ga or As sites depending on the growth conditions. Due to time constraint, no experiment could be made on GaAs(111)A layers. It is unsure if the available dopants (Si and Be) can be used efficiently in this orientation.

GaAs (111)B solar cells

In this section we fabricate the full solar cell structures on GaAs(111)B. We define the structure and the epitaxy parameters. After the growth we characterize the surface morphology. The sample is processed into individual device to extract statistics of the efficiency or other relevant parameters. The full solar cell structure is shown in Figure 4.10. We use a relatively simple design which can be easily transferred to other substrates. The N-type GaAs base is 1 µm and the p-type GaAs emitter is 100 nm, similarly to Wang et al. [Wang, 2017]. The InGaP material often found in the literature for the window and back surface field layers is difficult to grow due to the challenging control of the In content and the subsequent lattice match with the other GaAs layers. Instead, we use AlGaAs layers for which the lattice match with GaAs is always guaranteed [Wojtczuk, 1990;Urabe, 2015]. A single Al x Ga 1-x As composition is used (x∼0.3) for the windows and the back-surface field layers.

N-and P-type dopants are Si and Be (see [Lee, 2014] and the previous section). The Table 4.1 summarizes the respective layers nominal thickness and carrier concentrations.

All layers are grown using similar growth condition: As 4 BEP 1.2•10 -5 Torr, As 4 :Ga BEP ratio ∼12 and 1 µm/h growth rate. The substrate temperature is 620-630°C unless specified otherwise.

Layer

Thickness (nm) Doping (cm

Optical and SEM characterization

Results

We first characterize the solar cell structure by SEM and optical microscopy. Figure 4.11 (a) shows bird view SEM image at 45°tilt angle. We observe an important roughness created by small pyramidal holes (100-400 nm wide) on all the surface. Figure 4.11 (b) shows the associated optical characterization. The pyramidal holes are barely resolved as small dark dots at the maximum magnification.

We also observe a few larger features scattered on the surface.

Discussion

The observed surface roughness with a large density of relatively small pyramidal holes suggests that the growth conditions are not adapted. This type of morphology has previously been observed in chapter 3 (section 3.3.2.2, page 108) and it is associated to low As flux growth conditions. Future optimizations (section 4.5) will investigate the effect of the As flux and temperature on the solar cells morphology.

The total layer thickness is marginally inferior to that expected. Still the variations are well contained (∼10% error) and should not impact significantly the device performance.

The difference in the SE contrast with the doping type in semiconductor layer is well known [Elliott, 2002;[START_REF] Sealy | Mechanism for secondary electron dopant contrast in the SEM[END_REF]. The SEM cross-section analysis is thus very interesting as it combines the precise location of the top and bottom AlGaAs layers and a simple visualization of the location of the PN junction.

The observed pyramidal holes on the surface suggests a lack in As during the growth. The differences in layer thickness are small with respect to the target structure. The cross-section SEM is quite informative and show details of the solar cell structure, including the positions of the AlGaAs layers and that of the PN junction.

Metallic contacts

After the MBE growth, we proceed with the fabrication of the device, starting with the deposition of the metallic contacts by optical lithography and the etching of mesa to separate the cells.

Contacts fabrication

Figure 4.12 shows the fabrication process of the solar cell. The initial layer structure after the MBE growth is shown in Figure 4.12 (a). The fabrication procedure consists in the contacting the two junction extremities (top and bottom) and in the etching of the side of the cells to obtain individual devices.

After a short rinsing in HCl:H 2 O 2 3:7 for 1 min the sample is dried for 2 min at 125 °C. The resist (AZ5214) is spinned on the sample (30 s at 4000 rotation/min). A soft bake of 1 min at 125 °C is used to drive off the excess of solvents. Then the sample is exposed for 5 s with an inverted mask, and bake for 1 min at 125 °C to reverse the exposed resist, exposed for 30 s without the mask and finally developed (45 s in developer AZ826), see Figure 4.12(b).

Before the contact deposition, the sample is cleaned with 1 min HCl:H 2 O 2 3:7 to deoxidized the GaAs top surface. We then deposit the top contact by e-beam deposition using a mixture of Ti/Au 20/200 nm (Figure 4.12 (c)).

The top contacts are then defined by lift off, using acetone combined to ultrasounds, Figure 4.12 (d).

The exposed GaAs contact layer between the top Ti/Au contacts is etched for 70 s in (C 6 H 8 O 7 /H 2 O 1 gr/ml):H 2 O 2 5:1 (etching rate of ∼300 nm/min), see Figure 4.12 (e), to selectively remove and stop the chemical attack on the AlGaAs window layer. A second optical lithography step is used to etch the mesa, in which we protect the cells with resin before wet etching, Figure 4.12 (f). The sequence is: sample dry 2 min at 125 °C, spin of the AZ5214, soft bake at 125 °C for 1 min, insulation with UV for 15-18 s with the mask, development with MIF826 for 75 s, bake at 125 °C for 1-2 min. The exposed surface of the sample are then wet etched (unselectively) with H 3 PO 4 : H 2 O 2 : H 2 O 1:3:10 for ∼10 min. The etching rate is ∼100 nm/min. The final structure of the fabricated cells cell is shown in Figure 4.12 (g).

The top surface sample is then protected with resist and the sample is placed upside down in the e-beam deposition machine. We first deoxidize the GaAs bottom surface in HCl:H 2 O 2 3:7 for 1 min, then we deposit a stack of Ni/Au/Ge/Au/Ni/Au of 4/10/60/110/10/100 nm (Figure 4.12 (h)). The sample is then cleaned in acetone to remove the resist. The specific sequence of Au/Ge/Au 10/60/110 nm will form an AuGe eutectic if annealed. Figure 4.13 shows optical characterization of cells after the fabrication process. We will mainly focus on the largest cell which are ∼0.96 nm large in diameter. (internal active area ∼5.6•10 -3 cm 2 due to shadowing of the contacts) one another using linear fits of the measured current-voltage curve (I-V).

Results

In the case of top GaAs contact, we observe an ohmic contact, before and after the annealing, Figure 4.15 (a) and (b). The full TLM measurement are summarized Figure 4.15 (c). with the measured resistance GaAs p-type contact as a function of the distance between the contacts. The TLM measurements can simply differentiate the layer and contact resistivity as follows:

(4.1)

R T = R s W L + 2 • R c
with R T total resistance measured, L the length between the pads, W the width of the patterns (∼195 µm) and R s the semiconductor resistance. R c , the contact resistance it is half of the y axis intercept in the graph in Figure 4.15 (R T (L=0) /2).

The contact resistivity is defined as:

(4.2) ρ c = R s • W 2
The semiconductor resistance R s is known from the graph: R s = α • W , with α the slope of the graph in Figure 4.15 (c).

The top contact on the p-type Ga shows a very low contact resistivity of ∼ 3 mΩ • cm 2 before and after the annealing. The annealing may actually degrade the top Ti/Au contact characteristics and it is not required for this contact.

Bottom contact on n-GaAs

We repeat the same experiment on the back GaAs n-type contact. The full TLM measurement does not show the expected linear relation. However the resistance measured between two close pads is 1-2 , which place upper bound on the contact resistance.

Results

The formation of the Au-Ge eutectic is critical to obtain a good GaAs n-type back contacts of the solar cell.

Bottom contact on n-Ge

We repeat the same experiment as before but on the back Ge n-type contact.

Results

We observe a strong rectifying behavior with a large associated resistance before annealing (a) and a pure ohmic behavior after (b). The resistance drops significantly (several orders or magnitude) with the annealing procedure and the formation of the Au-Ge eutectic. The full TLM measurement result in a contact resistivity of ∼1.5 m • cm 2 .

As for the GaAs n-type back contact, the annealing of the Au/Ge metals is critical to obtain an ohmic contact at the back of the Ge (111) n-type wafer.

We are able to obtain good ohmic contact for the top and bottom contact. The top contact on the p-type GaAs does not require any further treatment while it is critical to anneal the bottom contact on the n-type GaAs (111) and n-type Ge (111) substrates.

Solar Cell characterization

To characterize the solar cell performance, we compare the current-voltage (IV) characteristic of the full device in the dark and under simulated solar light exposure. We also measure the External Quantum Efficiency (EQE) using a spectrometer to obtain current response of the solar cell when exposed to a mono-chromatic light source.

IV curve

Description of the experiment We first measure the current voltage (IV) characteristic of the solar cell. We use a Solar Simulators (Oriel LCS-100) and a Source Meter (Keithley's Series 2400 Source Measure Unit (SMU)). The standard illumination reproduced by the instrument is the AM 1.5G1 The solar simulator is made by xenon lamp with many filters to adapt the emitted spectra. Through a projection system, an uniform light beam is created on an area of about 10×10 cm.

The IV characteristic is measured by the SMU using the 4-probe configuration to avoid the parasitic resistance of the wires, under dark or light conditions. We only characterize the large cells shown in Figure 4.13 p. 166 (external diameter 0.96 mm). The current density is obtained dividing the measured current by the exposed active surface of the cell (A=0.0056 cm 2 ) In Figure 4.18 (b), we show the power density as a function of the voltage, computed from the previous I-V curve. We identify the maximum power density, P max , of the cell. From this result, we also extract the Fill Factor (FF) parameter which can be calculated as:

(4.3) F F = P max I sc • V oc
We can also calculate the cell Efficiency (Eff) in % with the formula:

(4.4) Ef f = P max P inc = F F • I sc • V oc P inc
The incident power on the cell is the incident power density of 100 mW/cm 2 multiplied by the cell surface area A.

The calculated parameters are resumed in

Discussion

We obtained an efficiency of 12.2% for the best cells but the average efficiently on the sample is about 9.2%. In details, the short circuit current is constant over the different cells but the V oc and thus the fill-factor and efficiency vary. The variations of the V oc indicates that the GaAs layer may present defects inside the layer and that the growth conditions are not optimized, which is also visible on the sample surface (Figure 4.11 p. 163).

We obtained an average efficiency of ∼9.2% for a GaAs solar cells (best cells 12.2%) on GaAs(111)B with no anti reflection coating. We show a reduction of fill-factor and V oc from cell to cell, possibly due to poorly adapted growth conditions.

External quantum efficiency

Description of the experiment We now characterize the EQE (external quantum efficiency) also said IPCE (Incident Photon-to-electron Conversion Efficiency). This measurement computes the ratio between the generated and impinging powers at each wavelength between 350 nm to 900 nm (step size is 5 nm).

We use here a tungsten filament, and a chopper, associated to a diffraction grating and a mirror. The ICPE is measured using by locking technique, by comparing the injected and output power of the solar cell.

Results

The EQE response of the best cell (12.2% efficiency) is shown in Figure 4.19. We observe a maximum at ∼600-700 nm. Discussion For low energy photon, the absorption is limited to wavelengths below the band gap of the GaAs of 1.424 eV (λ=871 nm). The associated GaAs absorption edge is clearly visible just below 900 nm. At low wavelengths the response is possibly limited by the surface recombination and the reflection coefficient.

In the previous I-V we measured a rather low current density of 15.6 mA/cm 2 . The absorption of the AM1.5G photons into GaAs and would yield a total current density of 28.6 mA/cm 2 for a 1 µm thick GaAs absorber and 31.1 mA/cm 2 for a 2 µm thick GaAs. This is calculated for a single light path, without considering any reflection. Theoretically, there is little improvement for absorber thicker than 1 µm. However our 1 µm thick GaAs emitter is only performing at 60% and may benefit from the additional thickness.

Conclusion

In this section we were able to fabricate a GaAs/AlGaAs solar cell starting from a commercial monocrystalline GaAs (111)B substrate. After the calibration of the n-type doping (Si) and the p-type doping(B), the full solar cell structure was grown and processed to obtain individual devices. The characteristics of top and bottom contact layers were investigated on test structure to obtain ohmic behavior. The final device show the typical features of a solar cell, with clear difference between light and dark I-V characteristics. The average efficiency is 9.6% and the best cell reach 12.2% (V oc = 0.97 V, J sc = 15.6 mA, FF = 81%).

Possible improvement routes are the increase of the n-type GaAs base layer and adjustment of the growth conditions.

GaAs (111)B solar cell optimization

Improvement of the current density

We first try to increase the n-type GaAs absorber from ∼1µm to ∼2µm, to increase the final current density. In an attempt to decrease the surface roughness, we also adjust the As flux so that the As 4 :Ga BEP ratio increases from 12:1 to 17:1.

Surface roughness and cross-section

We characterize the surface morphology with SEM and optical microscopy (Figure 4.20 (a,b)).

Results

The surface presents pyramidal holes at the SEM scale, Figure 4.20 (a). At Nomarski scale we observe conical islands 50-100 µm wide, Figure 4.20. Due to the doping contrast, we can verify using cross-sectional SEM the higher thickness of the absorber layer. Due to the increase in As 4 :Ga flux ratio the surface shows less pyramidal holes but additional As flux is still required to obtain a hole free surface.

IV characterization

We characterize the cells by IV characterization with the previously described solar light simulator.

Results

The obtained parameters from the IV characterization are reported in Table 4.3. We report the fill-factor, the short circuit current, the open circuit voltage and the calculated efficiency for the best cell and for the average over 11 cells of the same sample. IV curve in light and dark condition of the best cell is shown in Figure 4.21 with comparison with the best cell of previous section (section 4.4.3.1, p. 170). Discussion Compared to the first solar cell (see section 4.4.3.1), the current version with a thicker absorber allows to increase the current density from ∼15.6 mA cm 2 to ∼17.5 mA/cm 2 (best to best cell) and from ∼15.6 mA/cm 2 to ∼16.8 mA/cm 2 on average. However, this increase in current is balanced by a drop in the open circuit voltage of ∼0.1 V compared to the previous sample. This results in a lower overall efficiency from 12.2% to 10.5% (best to best cell) and 9.2% to 8.2% (average). There is thus no gain in increasing the thickness of the absorber layer over 1 µm.

A (cm 2 ) J sc (mA/cm 2 ) V oc (V) FF (%)
The drop in V oc with increasing thickness can be related to the poor material quality, despite the increased As flux and V:III ratio used in this experiment. This hints that the main limitation of the current, is caused by the poor material quality, which lead to numerous parasitic recombinations of the carriers.

In the next experiment we return to a thinner ∼1 µm GaAs absorber layer, with improved growth conditions. The higher thickness of the GaAs absorber (2 µm instead of 1 µm) allows us to increase the J sc of ∼2 mA. However, we lose in V oc and fill-factor, which degrade the overall efficiency of the cell.

Higher growth temperature and As flux

We try to improve on the initial solar cell (∼1 µm thick absorber), section 4.4.1.1 p. 162, using an even higher As 4 :Ga ratio of 27:1. Such a high value is obtained by both increasing the As 4 flux to its maximum value and decreasing the Ga flux from ∼9.5•10 -7 Torr to ∼6•10 -7 Torr. We also increase the growth from 620-630 °C to 640 °C. To correct for the growth rate decrease with the temperature, we slightly increase the growth duration for each layer.

Surface characterization

We first characterize the surface morphology with SEM and Nomarski (

Discussion

The combined use of higher growth temperature (640 °C) and higher As 4 :Ga ratio (27:1) succeeded in smoothing the sample surface, both at the large scale (optical images) and at the sub-micron scale (SEM). Yet, we still have to investigate if this improved morphology is representative of a better material quality for the solar cell.

High growth temperature (640 °C) and high As 4 :Ga BEP flux ratio (27:1) significantly improve the surface morphology of the topmost layer of the GaAs (111)B solar cell.

IV characterization

We then process the sample to obtain arrays of individual solar cells. The collected IV characteristic and resulting cells efficiency are shown The IV curves in light and dark condition of the best cell are shown in Figure 4.23 with the initial reference cell from section 4.4.3.1, p. 170. 3 (143) in light and dark condition with comparison of the IV curve of the best solar cell of section 4. 4.3.1, p. 170 Discussion We recover the higher V oc (0.96 V) of the cell obtained on the first iteration (4.4.3.1 p. 170). Coupled to slightly higher short circuit current, we obtain a small improvement of efficiency from 12.2% to 12.9% for the best devices. The main improvement is observed by looking at the average device, which is significantly improved (from 9.2% to 12.5% efficiency), with average characteristics very close to that of the best cell. This indicates that the sample is more homogeneous, which we attribute to the change in the substrate temperature and the V:III ratio used during growth.

Increasing the growth temperature (640 °C) and the As 4 :Ga BEP ratio (27:1) result in a smoother morphology of the layer, which is associated to a significant increase of the average device performance (from 9.2% to 12.5%).

Anti reflection coating

Description of the experiment The final optimization of this work will consist in adapting the optical index to minimize the light reflection of the top window layer. To this end, we deposit a anti-reflection coating (ARC) made of a single Si 3 N 4 homogeneous layer using a sputtering technique. The refractive index of the material is n∼2.0 and we used a thickness of 70 nm, adapted from existing optimization in the literature on GaAs solar cells [Wang, 2017].

We use here the best available solar cell, which is presented in section 4.5.2, and we compare their performance with and without ARC.

Results Figure 4.24 (a) shows the IV curves in light and dark conditions with and without the ARC. We observe that the ARC has the effect of shifting the light curve toward higher current density. A slightly lower open circuit voltage is observed. IV parameters are resumed in Table 4.5. Figure 4.24 (b) shows the normalized reflectivity (1-R) and the external quantum efficiency (EQE) with and without the ARC. We observe that ∼30% of the light was reflected before depositing the ARC. After deposition the reflected light reduces almost to zero at ∼550-600 nm. As a result, the EQE increase from ∼0.6 to ∼0.9 at 600-650 nm.

A (cm 2 ) J sc (mA/cm 2 ) V oc (V) FF Efficiency (%) 0.0056 21.6 0.92 80 15.9

Discussion

The most important reduction in the reflectivity appears in the range of ∼550-650 nm, which matches the highest improvement in the measured EQE. This confirms that the present ARC thickness and optical index are in the correct range, which may be improved by further dedicated modeling.

The reduction of the reflectivity directly translates as an increase of the photon inside the cells and a significant improvement of the short circuit current from 16.3 to 21.6 mA/cm 2 . However the deposition of the ARC is also results in a degradation of the V oc from 0.97 to 0.92 V. A possible explanation for the V oc decrease is that the plasma exposure used during the sputtering process damages the top layer of the cell. Overall the final efficiency increases from 12.5 to 15.9%, but we could have expected a larger gain in absence of the V oc decrease.

Our final device on monocrystalline GaAs (111)B wafer only reaches about 16% efficiency with antireflection coating, which is far from the 28.8% record efficiency on GaAs(100) substrates [Green, 2018]. It is actually comparable to that obtained by Wang et al. (16 %) using defective Ge buffer on Si (100) substrates [Wang, 2017] but slightly lower than the 18.2% obtained by venkatasubramanian et al. [Venkatasubramanian, 1996] on a large-grain poly-Ge substrate.

Our GaAs(111)B solar cell is thus not a record device, but it shows robust performance all across the sample area. However more work is needed to bring the efficiency of our exotic (111)B device to the level of those grown on standard GaAs (100) substrates.

Our final optimization using antireflection coating leads to increase efficiency to 16% for our AlGaAs/GaAs/AlGaAs solar cells on GaAs (111)B wafer. This value is far from the record cell of the literature, but it robust enough to evaluate the impact of other substrates on the device performance.

Conclusion

In this section we optimized an AlGaAs/GaAs/AlGaAs solar cell on GaAs(111)B monocrystalline substrates. After a first device fabrication, we identified a problem in the low short circuit current (J sc ) values. In a second cell sample, we attempted to increase the thickness of the GaAs absorber (from 1 to 2 µm), which did marginally increase the current at the cost of a decrease of the open-circuit voltage (V oc ) and fill factor (FF).

In a second iteration, we tried to improve the material quality by increasing the substrate temperature and the V:III ratio during the MBE growth. This approach improves the overall cell performance and results in homogeneous device characteristics across the sample.

We finally added an anti-reflection coating based on a single Si 3 N 4 layer to our best device. This resulted in a significant increase of the J sc and we measure of a final efficiency of ∼16% for our final device on monocrystalline GaAs(111)B wafer.

Although these performances are far from the record monocrystalline GaAs solar cells of the literature, they are robust enough to evaluate the effect of the substrate on the device characteristics, as we move from GaAs(111)B wafer to GaAs(111)A and Ge(111) wafers and later to MIC-Ge pseudosubstrates. On the Ge (111) sample grown at the same temperature, Figure 4.25 (c), we observe similar twinned area but larger pyramids. The twinned area appears as groups of small pyramids on the relatively smoother surface of the sample grown on Ge (111) at lower temperature, Figure 4.25 (d).

The top surface of the full solar cell structure grown on monocrystalline GaAs(111)A and Ge (111), show increase surface roughness compared to the reference GaAs (111)B reference, as well a twinned domains.

Considering the details of the fabrication procedure, see Figure 4.12 section 4.4.2.1 p. 164, in which the mesa etching occurs after the metal deposition, it is not possible to contact directly the substrate using the metal layers. This rules out basic issue during the fabrication procedure.

The previous results on the GaAs MBE on Ge(111) have yield GaAs(111)A epilayers (section 3.4.3.5 chapter 3 p. 138). Therefore the structures grown on GaAs(111)A wafer and Ge( 111) are (111)A oriented and only differ by the type of substrate. From section 4.3.1.2, p. 157, we know that the (111)A orientation is known to show severe compensation of the Si impurities, so that the carrier density of the Si-doped layer in the solar cell structure may be very low or even p-type.

However, the substrates are still n-type. Even if there is no PN junction in the solar cell structure, any p-type GaAs epilayer would still create a PN junction between the substrate and the epilayer. Therefore, either the Be impurities failed to produce a p-type doping the upper part of the GaAs solar cell or there exist another problem in our structure.

An alternative is the formation of a tunnel junction between highly doped p-type epilayers and the underlying n-type substrate, Ge (111) or GaAs (111)A. However, such a situation is rather unlikely since tunnel junction usually requires a fine tuning of the carrier densities and we do not observe any features related to tunnel junction in the IV characteristics, such as negative resistance [Esaki, 1966].

Solar cell structures on GaAs(111)A and Ge(111) substrates all show a severe short circuit behavior of undeterminate origin, which completely overcomes the device function.

MIC-Ge pseudosubstrates

Despite the severe problem identified above on monocrystalline Ge(111) wafers, we attempt to grow the same solar cell structure on GaAs(111)B. The primary goal of this experiment is to assess the morphology of the final layer, using a realistic solar cell structure and to identify MIC-related difficulties in the solar cell processing .

Sample fabrication

To avoid the selectivity problem with respect to silica for thick layer, see section 3.5.1 p. 144, we use here continuous large grain MIC-Ge on silica instead of patterned layer.

The MIC-Ge sample consist of a Ge/Al/GeOx/Ge (1/18/1.1/21 nm thick) on a silica layer annealed 10 h at 270 °C, following previous optimization (see section 2.3, chapter 2, p. 2.3).

Nomarski of the MIC-Ge substrate after etching is shown in Figure 4.27.

The MIC-Ge surface in then cleaned with H 2 O 2 (30 s) and HF (30 s) before the growth to remove any remaining Ge oxide. As we use a non-conductive silica layer, additional steps are required to deposit the bottom contact on the edge of each individual device. This is done by increasing the mesa etch time to expose the n-type base layer and to deposite the metal contact on the exposed area.

Characterization

Surface morphology We first characterize the surface of the cells by SEM and Nomarski microscopy, see Figure 4.28. We observe a very high surface roughness which can reach ∼750 nm, i.e. more than half the active structure.

Electrical characterization

Figure 4.29 show the I-V characteristic of a device fabricated on the rough surface of the layer. Under light, we observe a rather resistive behavior with no rectifying component in the studied voltage range. Under dark, the sample is very resistive.

Discussion

The peak to peak roughness of the layer is comparable to half the thickness of the active structure. This casts significant doubts on the contact deposition and the processing on the final device. It is very likely that we do not correctly contact the "bottom" and "top" layer of the solar cell.

In this disrupted geometry, the I-V characterization makes little sense from the device perspective. It is only used here as a simple comparison between the different substrates. Despite the uneven geometry, which offers many possibilities of short-circuit path and the short-circuit behavior observed on reference Ge (111) substrates (section 4.6.1), the devices fabricated on MIC-Ge pseudo-substrates are very resistive. This indicates that the surface roughness is also associated to large defect densities in the layers, possibly in the form of numerous grain boundaries, which significantly reduce the free carrier density. Compared to monocrystalline wafers, the growth of the full solar cell structure on MIC-Ge pseudo-substrates leads to very rough layers. The fabricated devices are all resistive, despite the many possible short-circuit path, which actually demonstrates that the GaAs layers grown on the MIC-Ge are very defective.

Conclusion on solar cells

In this chapter, we first optimize the fabrication of a simple AlGaAs/GaAs/AlGaAs solar cells in the (111)B crystalline orientation. We successfully fabricate solar cell device, starting from the MBE growth on the GaAs(111)B wafer, moving to the metallic contact deposition and mesa etching to fabricate the final arrays of individual devices. By optimizing the substrate temperature and the V:III ratio using during the growth we are able to increase the efficiency from 9.2% to 12.5% on average device, and then to 15.9% by adding a simple anti-reflection coating. The final devices on GaAs (111)B show robust performance but compare poorly to record GaAs solar cells obtained on GaAs (100) substrates.

After this optimization on GaAs(111)B, we fabricate the same solar cell on monocrystalline GaAs(111)A and Ge(111) wafers and MIC-Ge pseudo-substrates. In all cases, we fail to obtain working devices. On Ge(111) and GaAs(111)A, we obtain (111)A oriented devices which all present a short-circuit. Alternatively the layer growth on MIC-Ge pseudo substrate is so irregular that the roughness of the final structure does not permit to correctly contact the required layers. A tentative electrical characterization reveals very resistive layer, which indicates that the III-V material grown on the MIC-Ge is of poor quality.

In conclusion, we identify two severe problems in the fabrication of the solar cell on MIC-Ge pseudo substrate, first related to the poor crystal quality of the III-V material and second to the (111)A crystal orientation. Even if future optimization of the MIC-Ge template could lead to an improved crystal quality in the III-V, the growth of (111)A solar cell will likely require other doping impurities than Si and Be used in this work.

General conclusions

In the introduction, we divided the work towards the growth of GaAs solar cells on low-cost Ge pseudo substrates in three main challenges:

1. The fabrication on a silica substrate of large isolated monocrystalline Ge grains with a strong crystal texture, using Metal Induced Crystallization (MIC).

2. The selective growth of GaAs (111) layers by MBE on the Ge pseudo-substrates obtained above.

3. The MBE growth of a full GaAs-based solar cell structures and the fabrication individual devices on these Ge pseudo-substrates.

After the tedious optimization of the main experimental parameters (layer thicknesses, interlayer, underlayer, . . . ) of importance for the MIC of amorphous Ge using aluminium (otherwise called Aluminium Induced Crystallization, or AIC), we reached a point where further progress requires too many samples if only ex situ characterizations were considered. Instead, we assembled an in situ optical microscope to determine the precise chronology of the processes at play during the AIC of Ge and to measure their respective kinetics. While the AIC process is only associated to a single mechanism in the literature (ALILE), we identified two distinct processes, that we called MIC-1 and MIC-2. We then fully determined the morphology and orientation of the crystals produced by each process, using a combination of ex situ (SEM, TEM, EBSD) and in situ (optical imaging) characterizations. The better understanding of each MIC process was critical to let us limit the spatial extension of the defective dendritic MIC-1 crystals, and favor the MIC-2 process, which creates the large (111) oriented Ge grains. The use of lithography to obtain separate grains was partially successful. Still, we obtained 20 nm thick Ge platelets which extend laterally over several tens of microns and show a strong Ge(111) texture, on silica substrates.

Among the various possible strategies to grow GaAs on MIC-Ge pseudo-substrates, we investigated the growth of a Ge buffer and the direct growth of GaAs on the Ge pseudo-substrates. The later was found to be the most promising route, with the fabrication of relatively smooth and uniform GaAs layers, at least at the initial stages of the growth. The specific growth conditions of GaAs layers on the (111)-textured Ge platelets were optimized from that of reference GaAs (111) layers grown on monocrystalline GaAs(111)B and Ge(111) wafers. Contrary to the standard (100) orientation, we find that the growth of GaAs (111) oriented layers on Ge substrates does not produce anti-phase domains (with bonds between like atoms) but only twin defects, which density can be almost reduced to nil if the appropriate growth conditions are used. We observe and characterized the growth on Ge(111) surfaces of GaAs layers with the sole (111)A orientation, which is not the polarity expected from the literature. We also confirm the correct epitaxy of a GaAs layer on the (111)textured MIC-Ge crystals. Despite growth selectivity issues and the high roughness of the thick GaAs layers grown on the MIC-Ge pseudo-substrates, we proceeded to the growth of solar cell structures on Ge(111) surfaces, using either mono-crystalline wafers or MIC templates.

In the last stage of this work, we fabricated and optimized a simple GaAs solar cell structure on GaAs(111)B monocrystalline substrates. The average performance of the fabricated device is relatively modest, just below 16% efficiency, far from the record devices fabricated on GaAs(100) substrates (28%). However the growth and the fabrication process are robust enough to evaluate the impact of the different substrate on the device performance. The move from GaAs(111)B to GaAs(111)A and Ge(111) monocrystalline wafers is actually very difficult and only short-circuited devices are obtained, despite the similar fabrication procedure. This shows that the polarity of the layer is a critical parameter when optimizing layer doping or device fabrication. An attempt at fabricating the full solar cell structure on a MIC-Ge pseudo-substrate confirms the previous selectivity and roughness issues. It also points out the poor material quality of the GaAs epitaxially grown on the MIC-Ge pseudo substrates, since only very resistive layers are obtained.

In principle, the fabrication of MIC-Ge patches on silica appears as a promising alternative route for the fabrication of low cost GaAs solar cells. However, we find that the growth of smooth and good quality GaAs layers on our MIC-Ge pseudosubstrates is extremely challenging. Even if the GaAs polarity issue are solved to obtain (111)B layers on Ge (111), to avoid possible n-type doping problems, the high roughness of all the thick GaAs layers grown on MIC-Ge templates has remained impossible to overcome and points at intrinsic limitations of the crystal quality in MIC-Ge layers.

A.1. Powder X-ray diffraction

(A.2) d(h, k, l) = a √ h 2 + k 2 + l 2
• a is the unit cell parameter (a = 5.657 35 Å for Ge)

• (hkl) are the Miller indexes of the considered plane.

• θ is the scattering angle, half of the angle between source and detector (2θ).

In the conventional θ-2θ configuration (Bragg-Bentano), the incident and diffracted beam are inclined symmetrically to the surface, (ω = θ = 2θ/2) so that we probe the planes parallel to the surface.

In the case of monocrystalline sample, the inclination of the source to the detector (ω) is a critical parameter but it is not relevant in a powder sample as some crystallites will always be in diffracting conditions. The scanning speed of angle ω, between the incident beam and the sample surface, is set to half that of angle 2θ, between the impinging beam and the diffracted beam, to maintain the geometry. The scan can be made by rotating the sample and the detector (θ-2θ) or the detector and the source (θ-θ).

A.1.2 Diffraction of diamond cubic crystals and structure factor

For each planes, the diffracted intensity can be simply written as the multiplication of different terms, equation A.3. F is the structure factor; M is the multiplicity factor; L is the Lorenz factor; P is the polarization factor and T is the temperature factor.

(A.3)

I = F • M • L • P • T
In a first approximation, the structure factor F gives the main shape of signal (i.e. most intense reflection, forbidden reflections, etc...) and the other parameters modulate this value. For Ge diamond cubic crystals, the structure factor is obtained by the following equation:

(A.4) F = 2f 2
Ge 1 + (-1) h+k + (-1) k+l + (-1) h+l fcc term

• 1 + cos 2π(h + k + l) 4

diamond cubic additional term

The first term, f Ge , is the scattering factor of individual Ge atoms, which can be calculated from the CromerMann coefficients. The two others are the Face Centered Cubic (FCC) and the diamond interference terms. The first relates to the coherent diffraction of each FCC sub-lattice and the second is the interference between the two FCC sub-lattices shifted by a/4.

In practice, the FCC term only allows diffraction for hkl indexes all odd or all even, and the additional diamond cubic factor further restricts the "all even" reflections to that with h + k + l multiple of four.

A.1.3 Multiplicity in cubic crystals

The multiplicity is the number of planes which give the same diffraction angles. In a cubic system, the planes are equivalent by inversion h → -h (for the non-zero terms) and by the permutation of the different indexes.

We first calculate the number of equivalent term depending on the number of non zero terms: n°equivalent planes 0 0 x 2 0 x x 4 x x x 8

Then, we calculate the number of equivalent planes as a function of the possible permutations: n°equivalent planes h k l 6 h h l 3 h h h 1

We now compute the total multiplicity for each family of reflection: hkl hhl hhh hk0 hh0 h00

Multiplicity 8×6=48 8×3=24 8×1=8 4×6=24 4×3=12 2×3=6

The multiplicity is an important correction in powder samples as it is a straightforward multiplication of the signal.

A.1.4 Application to Ge powder

A.2 Crystal texture of Ge(111)

We move here to the θ-2θ characterization of a powder sample with a partial (111) crystal texture.

A.2.1 Qualitative analysis

In a powder, each grain is randomly oriented ( For the Ge powder in the standard θ-2θ configuration, only the (111) and the (444) planes are allowed to diffract.

A.2.2 Lotgering factors

Instead of a pure texture, samples can have a partial crystalline texture. Lotgering introduced the parameter p as a measure of the degree of orientation in a specific direction [START_REF] Lotgering | Topotactical reactions with ferrimagnetic oxides having hexagonal crystal structures-I[END_REF].

The parameter p is defined as p Here, we adapt the Lotgering factor L

A.4 Theoretical value of Lotgering factor in GIXRD

We study now the theoretical diffraction intensities and Lotgering factor of fullytextured samples analyzed in the 2θχ configuration.

A.4.1 (111) crystal texture of Ge

The (111) crystal texture is defined as the crystalline phase with the [111] direction perpendicular to the surface of the sample. In this particular orientation, the Ge crystal has an hexagonal symmetry (Figure A.7). Only the {110} and {211} planes are orthogonal to the (111) plane. Due to the hexagonal symmetry the multiplicity of each set of reflection is 6. It is not the same as in a fully random sample (Table A.1).

This specific multiplicity corrects the theoretical diffracted intensities. From the reference powder data, we compute the GIXRD values by renormalizing between the different multiplicities.

(A.8)

I GIXRD = I XRD • M GIXRD M XRD
For an easier comparison the diffracted intensities are rescaled to 57 for the 220 reflection so that the values between randomly oriented and full (111)-texture can be compared directly. The intensities are summarized in Table A The (100) texture presents a quadratic symmetry in the plane. Only the {110}, {100} and {210} plane families are orthogonal to the (100) plane. All planes perpendicular to the [100] take the (0hk) form due to the scalar product of the two vectors:

   1 0 0    •    0 h k    = 0.
For the multiplicity, we perform the same analysis than in the 3D case (see section A.1.3)

A.4.3 Numerical variations between 100 and 111 Lotgering factors

Table A.4 summarizes the intensities and multiplicities for a Ge powder, a full (111)and (100)-texture in the GIXRD configuration. The powder values are the same as in the standard XRD configuration.

We calculated the theoretical Lotgering L A possibility to improve the Lotgering factor would be to introduce weight when summing the reflections to minimize the contribution of the common terms.

A.5 Example of experimental result

We now consider the spectrogram of the sample "C" shown in chapter 2 (section 2.6.1), after wet-etching of Al, Figure A.9. We aim to compute the diffracted intensities, which are the areas of the peaks in the diffractogram. First we remove the background with a polynomial fit considering of the spectra outside the peaks. Secondly, we integrate each peak area separately.

From the experimental values and the reference powder data (table A.1) we obtain the following Lotgering parameters: As expected from the theoretical study, both the L (111) GIXRD and L (100) GIXRD show a high value. Having calculated both parameters let us know that the (111) texture is more likely than the (100) texture.

However the Lotgering analysis is very limited in the GIXRD geometry. To confirm the crystal texture, the mapping of the sample by EBSD is required. the parameter L as a measure of the degree of orientation, define as in formula ?? [START_REF] Lotgering | Topotactical reactions with ferrimagnetic oxides having hexagonal crystal structures-I[END_REF]Furushima, 2010].

A.6 Conclusion

The grazing incidence configuration is a convenient geometry for the analysis of thin crystalline layers. However, it is difficult to discriminate between (111) and (100) textures using Lotgering factors because the 220 reflection (highest intensity) is common to both textures. Alternative crystal structure characterization must be considered such as EBSD.

Annexe B

Résumé en français B.1 Introduction

De nombreux efforts ont été déployés pour développer des cellules solaires efficaces et réduire le coût de la production d'électricité photovoltaïque (PV).

Les principaux moteurs de ce développement sont l'augmentation des prix du pétrole, le réchauffement climatique, l'effet de serre et l'augmentation de la demande énergétique mondiale [Lewis, 2010].

Le GaAs est un bon candidat pour atteindre des rendements de conversion supérieurs à ceux atteints avec les cellules solaires à base de silicium. Ce dernier a une bande interdite indirecte de 1,12 eV tandis que GaAs a une bande interdite directe de 1,424 eV. En conséquence, une cellule en GaAs n'a besoin que de 2 à 4 µm d'épaisseur pour absorber l'ensemble des photons accessibles (avec energies superieur a la bande interdite) du spectre solaire au lieu de 100 à 200 µm pour le silicium. En outre, GaAs détient le record d'efficacité pour les cellules solaires à simple jonction (28,8 %) [Green, 2018].

Pour compenser la différence de prix d'un facteur 10 entre les substrats Si et GaAs, l'idée principale est de recycler le coûteux substrat GaAs et de le ré-utiliser pour plusieurs cycles de fabrication. Cependant, cette technologie de recyclage du substrat reste limitée dans le domaine du photovoltaïque et n'a pas encore fait ses preuves.

Dans ce travail, nous étudions une autre façon de fabriquer des substrats à faible coût pour la croissance épitaxiale de GaAs. Ces substrats sont composés d'une fine couche constituée d'ilots de Ge cristallin uniformément orientés. Ces pseudosubstrats sont fabriqués par Cristallisation Induite par les Métaux (MIC), à partir d'une couche de germanium amorphe déposée sur un substrat de verre. Bien que le procédé MIC fonctionne également avec d'autres matériaux, tel que le silicium, nous avons choisi le germanium car le paramètre de maille de ce dernier est proche de celui de GaAs. L'objectif final de ce travail est de réaliser une cellule solaire sur un support en verre en utilisant une couche cristalline obtenue par MIC (appelée MIC-Ge dans la suite) comme pseudo-substrat. Ce travail est divisé en trois parties qui correspondent à trois défis différents (Figure B.1). • Nous examinons d'abord le processus MIC appliqué au Ge amorphe. En particulier, nous développons un nouveau microscope optique pour suivre in situ la cinétique des processus de cristallisation.

MIC-Ge

• Ensuite, nous étudions l'épitaxie de couches de GaAs par épitaxie par jets molécuaires (MBE) sur des substrats monocristallins et sur des pseudo-substrats MIC-Ge. Nous caractérisons des défauts tel que les macles, spécifiques à l'orientation (111) des substrats.

• Pour finir, nous fabriquons et caractérisons des dispositifs à cellules solaires GaAs sur différents substrats : substrats monocristallins GaAs(111)B et Ge(111) et pseudo-substrats MIC-Ge sur silice.

B.2 Cristallisation induite par les métaux

La "cristallisation induite par les métaux" (MIC) est la cristallisation d'un matériau amorphe en contact avec un métal. La température de recuit est toujours inférieure à la température eutectique du mélange métal-semi-conducteur, de sorte que les matériaux restent en phase solide.

La technique MIC permet d'obtenir des films minces polycristallins à partir de couches amorphes de silicium ou germanium.

La cristallisation induite par l'aluminium (AIC) est également appelée ALILE (Aluminium Induced Layer Exchange), car les couches de Al et de Ge (Ge amorphe) échangent leur position grain par grain pendant la cristallisation du Ge [Nast, 2000b].

Nous montrons ici que des couches minces de Ge (d'épaisseur 20 nm) peuvent être cristallisées à condition d'introduire une couche intermédiaire de (Ge,Al)O x entre les couches d'Al et de Ge amorphe et une sous-couche (∼1 nm de Ge sous la couche d'Al), toutes deux soigneusement ajustés. Cependant, les conditions optimales de recuit dépendent fortement de la structure de l'échantillon.

B.2.1 Microscopie in situ

B.2.2 Caractérisations ex situ

Nous étudions les couches créées par MIC-1 et MIC-2, avant et après le recuit, afin de déterminer leur mécanismes de croissance respectifs.

La microscopie électronique en transmission (TEM) permet de caractériser structurellement et chimiquement les zones avant et après cristallisation.

Ceci nous permet de démontrer que le processus MIC-2 est le mécanisme standard de cristallisation par échange de couches induit par l'aluminium (ALILE), tel que décrit dans la littérature [Nast, 2000b ;Kurosawa, 2012]. En revanche, les grandes dendrites formées au stade MIC-1 sont constituées d'une par une double couche de Ge et apparaissent avant le processus ALILE (MIC-2).

Nous étudions ensuite l'orientation cristalline locale de chaque grain par diffraction des électron rétrodiffusés (EBSD). Nous observons que les cristaux de Ge forment des îlots isolés (d'environ 30 à 40 µm de diamètre), présenstant une texturation dominante delon la direction (111) ( La combinaison de la microscopie in situ et de l'EBSD nous permet d'affirmer que c'est bien le processus MIC-2/ALILE qui s'étend à la quasi-totalité de la surface de l'échantillon en formant des cristallites de Ge(111) orientés. A l'inverse le processus MIC-1 est associé à la formation des dendrites centrales, contenant à la fois des cristaux de Ge orientés (100) et (111).

Pour améliorer la qualité des pseudo-substrats MIC-Ge, il faut donc minimiser le processus MIC-1 en faveur du processus MIC-2.

B.3.1 Epitaxie de GaAs sur substrats de GaAs(111)B et Ge(111)

Sur les substrats de GaAs(111)B nous trouvons que la croissance optimale (en terme de planéité) nécessite une température du substrat de 630-640 °C et un flux d'As 4 spécifique.

Nous passons ensuite à l'étude expérimentale de la croissance de GaAs sur des substrats de Ge(111) monocristallins. Nous caractérisons la rugosité de surface du GaAs, après des croissances effectuées avec des flux d'As et des températures du substrat différents. Nous observons également par diffraction des rayons X (XRD) des macles dans la couche épitaxiale de GaAs. La densité des macles est très importante pour les flux d'arsenic les plus bas et les plus élevés.

En combinant la cathodoluminescence (CL) et la microscopie TEM, nous montrons que les domaines maclés sont séparés par des joints de grains présentant une grande densité de centres de recombinaison non-radiative. Seule la croissance de l'échantillon utilisant un flux d'arsenic intermédiaire permet d'obtenir un matériau de qualité, avec une faible densité de macles et de joints de grains. D'après la littérature, lorsque la surface de Ge(111) est exposé à l'As, elle se passive entièrement, sans présenter de liaisons pendantes [Bringans, 1985]. Cette surface, terminée par les atomes As, devrait conduire à la formation exclusive de couches de GaAs d'orientation (111)B [Koh, 1999 ;Koh, 2001]. Cependant, nos analyses chimiques locales en TEM montrent sans ambiguïté une polarité (111)A du GaAs épitaxié sur Ge (111). Cette différence importante par rapport à la littérature pourrait être due à des conditions de croissance différentes.Nous avons effectué une analyse STEM-EDX résolue atomiquement qui montre sans ambiguïté une polarité de GaAs(111)A dans nos échantillon sur Ge (111). La différence par rapport à la littérature pourrait s'expliquer par des conditions de croissance différentes.

B.3.2 Epitaxie de GaAs sur MIC-Ge

Il s'agit ici d'effectuer la croissance épitaxiale de GaAs sur les pseudo-substrats MIC-Ge, en tenant compte des résultat obtenus sur les substrats monocristallins (section B.3.1). Après optimisation, nous trouvons que la croissance directe de GaAs sur le pseudo-substrat MIC-Ge n'est que partiellement sélective par apport à la silice. En revanche, elle produit une couverture uniforme de GaAs sur les plaquettes de MIC-Ge. Cependant, plus la couche de GaAs sur MIC-Ge est épaisse, plus elle devient rugueuse (plus de 500 nm de dénivelé pour une épaisseur moyenne de 2 µm). De même, la sélectivité de croissance par rapport à la silice est perdue à partie de 360 nm d'épaisseur. Sur les pseudo-substrats MIC-Ge, la croissance de la cellule solaire complète créée un rugosité extrêmement importante (plusieurs centaines de nm), incompatible avec la prise de contact. Une mesure électrique préliminaire montre que le matériau est très isolant, ce qui implique que les couches de GaAs épitaxiées sur MIC-Ge contiennent de nombreux défauts.

B.5 Conclusion

Nous avons tout d'abord optimisé les paramètres de fabrication par le procédé MIC-Ge. Pour ceci, nous avons dû mettre en place une caractérisation optique in situ de la cristallisation pendant l'étape de recuit thermique. En combinant les caractérisations in situ et ex situ, nous avons observé et distingué deux sous-processus, baptisés : MIC-1 et MIC-2.

Tous deux sont des processus de cristallisation induite par un métal qui transforment la couche de Ge amorphe, mais les cristaux obtenus ont des structures différentes. Le processus MIC-1 crée des doubles couches de Ge où seulement la couche inférieure est cristallisée, avec une texture mixte ( 100 .9 -Schémas du processus MIC-1, qui crée une couche de Ge cristallin (c-Ge) en dessous de la couche initiale de Ge amorphe (a-Ge) [START_REF] Pelati | In Situ Optical Monitoring of New Pathways in the Metal-Induced Crystallization of Amorphous Ge[END_REF], comparé au processus classique MIC-2 (ALILE) qui procède par l'inversion des couches de Ge et Al [Nast, 2000b].

En ce qui concerne l'épitaxie de GaAs par jets moléculaires, nous avons d'abord étudié des couches de références sur substrats monocristallins de GaAs(111)B et Ge (111). Sans aucune couche tampon, nous avons montré que l'ajustement du rapport de flux As :Ga joue un rôle essentiel dans l'obtention d'une faible rugosité et pour éviter la formation de macles lors de la croissance de GaAs sur Ge(111). De manière assez surprenante, nous observons que les couches de GaAs adoptent une polarité (111)A sur Ge (111), alors que la littérature mentionne principalement des couches de polarité (111)B.

Figure 1 . 1 -

 11 Figure 1.1 -Graphical summary of the thesis: fabrication of MIC-Ge pseudo-substrates on silica (chapter 2), GaAs epitaxy on MIC-Ge (chapter 3) and processing of solar cells (chapter 4).

Figure 2 . 1 -

 21 Figure 2.1 -Ranges of crystallization temperature of Ge by Solid Phase Crystallization (SPC), Metal Induced Crystallization (MIC) and laser annealing.

Figure 2

 2 a-Si is deposited on the top of the Al layer.2. Si atoms diffuse from the a-Si into the metal. The Al grain boundaries acts as fast diffusion paths for the Si atoms. As shown in Figure2.2 (b), the Si nucleation occurs at the Al grain boundaries, when the concentration of Si

Figure 2 . 2 -

 22 Figure 2.2 -Schematics of the MIC steps: a-Si and Al are in contact, which allows Si to diffuse inside Al; Si nucleates inside the Al layer; c-Si grows laterally and pushes Al to the top position; the two layers have finally inverted their position (left). Cross sectional SEM micrograph at these successive MIC steps (extracted from[Nast, 2000b])(right)

Figure 2 . 3 -

 23 Figure 2.3 -Patterned substrate. Fraction of crystallized patches as a function of the lateral size of the patches, for different annealing times (a) and as a function of annealing time for different sizes (b).(c-e) Gaussian model for patch amorphization during the etching of the trenches between the patches (extracted from[START_REF] Cohin | Crystallization of Si Templates of Controlled Shape, Size, and Orientation : Toward Micro-and Nanosubstrates[END_REF]).

  Figure 2.4 (a) and Figure 2.4 (b) show optical images of the SiGe

Figure 2 . 4 -

 24 Figure 2.4 -Nomarski optical micrographs of a -Si 0.6 Ge 0.4 samples after annealing for 30 min at 650 °C (a) and 600 °C (b). Temperature below which MILC is possible as a function of the Ge fraction in a-(Si,Ge) samples (c). (extracted from[Kanno, 2006]).

Figure 2

 2 Figure 2.6 -(a) Phase diagram of the Al-Ge system (from[START_REF] Mcalister | [END_REF]); (b) Phase diagram of the Au-Ge system (from[Okamoto, 1984])

Figure 2 . 7 -

 27 Figure 2.7 -Left: schematic showing the diffusion of the Ge atoms into the Al layer through the AlO x interlayer and the resulting nucleation at the metal/substrate interface.Right: schematics of the expected Ge concentration inside the layers (extracted from[Toko, 2014b])

Figure 2 . 8 -

 28 Figure 2.8 -Electron backscatter diffraction map showing the crystalline orientation, of MIC samples (with Al and Ge thicknesses of 50 nm) as a function of annealing temperature and oxidation time of Al in air. Annealing time are: (a) 10 h,(b) 30 h, (c) 100 h, (d) 10 h, (e) 30 h, (f) 100 h, (g) 30 h, (h) 100 h, and (i) 400 h. (extracted from[Toko, 2012])

Figure 2 .

 2 Figure 2.9 -(111)-oriented fraction of Ge crystallized by AIC as a function of the Ge atomic content in (Si,Ge) material and of the Al layer thickness(50, 75, 100 nm). The (Si,Ge) layer has the same thickness as the Al layer (extracted from[Kurosawa, 2012])

Figure 2 . 10 -

 210 Figure 2.10 -Main steps of sample preparation and annealing in the Aluminium Induced Crystallization (AIC) of amorphous Ge (a-Ge) into crystalline Ge layer (c-Ge). A thin (1 nm) amorphousGe layer is used as an underlayer. The AlO x or GeO x interlayer is obtained by air oxidation of the Al layer (from[START_REF] Pelati | In Situ Optical Monitoring of New Pathways in the Metal-Induced Crystallization of Amorphous Ge[END_REF]).

Figure 2 .

 2 Figure 2.11 -Schematic of the PVD deposition of Al and amorphous Ge by electron beam evaporation.

Figure 2 .

 2 Figure 2.12 -The PLASSYS MEB 550 SL e-beam evaporation machine used at C2N for Al and Ge evaporation.

Figure 2 .

 2 Figure 2.13 -(a)Schematics of the three-zone AET oven. In green, the external thermocouple used for the calibration of the oven; (b): Plot of the measured temperature as function of the oven setpoint (c). Plot of the temperature offset as a function of the oven setpoint

Figure 2 . 14 -

 214 Figure 2.14 -Schematics of Nomarski DIC microscopy, with off-axis ray traces.

Figure 2 .

 2 Figure 2.15 -Optical images of a MIC sample obtained in bright field mode with standard optics (a) and with Nomarski differential interference contrast (b).

Figure 2 . 16 -

 216 Figure 2.16 -Nomarski optical images of the MIC-Ge samples, (see summary of properties in Table 2.1) after annealing and etching of the Al layer. The thickness common to Al and Ge layers increases from left to right, as indicated on top axis. Scale bar is 20 microns.

Figure 2 . 18 -

 218 Figure 2.18 -Nomarski optical images of samples after annealing and Al etching, as a function of the Ge thickness: ∼16 nm (#1), ∼20 nm (#2) and ∼24 nm (#3), for a fixed Al thickness of ∼20 nm. (a-c) Schematics of the cross-section of the samples listed in Table 2.3; (c-d) Nomarski top views; (f-h) zoomed views of (c-e).

Figure 2 .

 2 Figure 2.19 -Comparison of Nomarski optical images of samples with an AlO x interlayer (a), with an AlGeO x interlayer (b,c,d), and without any interlayer (e). Thickness of the as-deposited Ge layer (before oxidation): 1,1 nm (b), 1.5 nm (c) and 3 nm (d). For the details of these samples, seeTable 2.4.

Figure 2 . 20 -

 220 Figure 2.20 -Schematics of the time required to complete the crystallization as a function of the type of interlayer oxide.

Figure 2 . 21 -

 221 Figure 2.21 -Nomarski optical images of the MIC-Ge samples with: no Ge underlayer (a), 1 nm Ge underlayer (b) and 2 nm Ge underlayer (c).

Figure 2 . 22 -

 222 Figure 2.22 -Photograph and schematics of the oven hotplate fitted with optics.

Figure 2 .

 2 Figure 2.23 -In situ temperature monitoring during sample annealing for over than 60 h.

Figure 2 .Figure 2 . 25 -

 2225 Figure 2.24 -in situ optical characterization of MIC sample at 45 min (a), 60 min (b) and 90 min (c). Measured temperature as a function of the annealing time

Figure 2 . 26 -

 226 Figure 2.26 -Schematics of sample position inside the annealing oven and optical access.

Figure 2 .

 2 Figure 2.27 -In situ optical microscopy from the back of the sample (Al-side, sample 42D8) during the AIC annealing at 238 °C. (q) and (p) are zoomed views of the areas indicated by rectangles in (d) and (i), respectively. Scale bar is 40 µm. Black arrows indicate the MIC-1 nucleation and growth. Red arrows indicate the onset of the MIC-2 process. (l) scheme of the MIC-1 and the MIC-2.

Figure 2 .

 2 28 presents the in situ optical monitoring of the sample without the Ge underlayer.

Figure 2 .

 2 Figure 2.28 -In situ optical microscopy of a sample with no Ge underlayer annealed at 255 °C, observed from the back. Insert in (a) shows the sample structure. (a)-(d) Optical images at annealing times of 0.6 h, 0.7 h, 0.88 h and 2.45 h respectively. (e)-(g) Zoomed views of the area indicated by rectangles in (b)-(d). (h) Schematics of MIC-1 and MIC-2 areas with reference to zoomed image (f). Scale bars: (a)-(d): 20 µm, (e)-(g): 5 µm. Black arrows indicate the starting point of the MIC-1 dendrite. Red arrows indicate the onset of the MIC-2 process.

Figure 2 .Figure 2 . 30 -

 2230 Figure 2.29 -In situ image processing with ImageJ: (a) before processing; (b) after conversion to grayscale and image shift; (c) after image thresholding and lines along which dendrite lengths are measured.

Figure 2 .

 2 Figure 2.31 -Spatial extension of the MIC-2 processes as a function of time, using length measurements L1 and L2, (see inserts, scale bar is 40 µm).

Figure 2 . 32 -

 232 Figure 2.32 -Schematics of the sample position and optical observation inside annealing oven.

Figure 2 .

 2 Figure 2.33 -(a) Spatial extension of the MIC-1 (a) and MIC-2 (b) processes as a function of time viewed from the top. In (a), letters A, B, C refer to three different branches of MIC-1 dendrites (inserts show images taken at times indicated by dotted lines, scale bar is 40 µm). (b) Spatial extension of the MIC-2 process using either direct length measurements (L1, L2), or approximating the area covered by a disk. (see insert). Scale bar: 100 µm.

Figure

  Figure 2.33 (b) shows how the expansion length varies with time in the MIC-2 process. This process does not consist in individual dendrites but rather produces uniform light areas. Since these areas are somewhat irregular, we estimate the MIC-2 extension either by direct linear measurements (L1 and L2 in the figure) or as the radius of the disk having the same area as the MIC-2 zone. Both measurements give a rate of 2.5 nm/s over 10 h annealing at 237± 2 °C, is. As previously observed from the bottom, the MIC-2 areas uniformly expand and merge to cover the entire surface of the sample.

Figure 2 .

 2 Figure 2.34 -Nomarski optical image (a), SEM and EDXS images at 2 kV (b,c) and 7 kV (d,e) of a partially annealed sample (6 h at 260 °C) sample.

  An example of CASINO simulation of a Ge/Al/SiO 2 /Si stack with thicknesses 21/20/39 nm at 2 kV is shown in Figure 2.35 (d). Simulations of the measured Ge/Al ratio or Al/Ge ratio (full lines) for different incident electron energies (2 to 10 kV) are plotted together with the corresponding experimental results (dots) in Figure 2.35 (b) and (c). The simulated thicknesses and position for Al and Ge in each of the E0, E1, E2 and E3 zones are reported in Figure 2.35 (e).

Figure 2 . 35 -

 235 Figure 2.35 -Example of EDXS spectra after background subtraction and integration of the peaks (a). Ratios of peak intensities (Ge/Al (b) and Al/Ge (c)) for the four zones shown in Figure 2.34 (d). CASINO simulation for SEM acceleration voltages from 2 kV to 10 kV: (d) Example of CASINO simulation. (e) Summary of the structures best matching the EDXS measurements.

Figure 2 .

 2 36 (b)) shows a very clear chemical contrast. The MIC-1 dendrites (labeled M1), producing a high Ge signal and very little Al signal, are clearly distinguished from the relatively Al-rich MIC-2 zones (labeled M2).

Figure 2 .

 2 Figure 2.36 -Top view SEM image (a) and associated EDXS map at 5 kV (b) showing the Al (red) and Ge (light blue) signals of a partially annealed sample. The label R indicates the reference area, M1 the dendrite created by MIC-1 and M2 the surface created by MIC-2. The dashed white line indicates the FIB cut used to extract a TEM sample containing the interfaces R/M2 (pink rectangle) and M1/M2 (blue rectangle). (c-e) TEM-EDXS maps of the Ge (light blue), Al (red) and O (green) signals from R/M2. (f-h) TEM-EDXS map from M1/M2. Scale bars in (c-h) are 100 nm.

  (a) and 2.36 (b).

Figure 2 . 37 -

 237 Figure 2.37 -Schematics of the new MIC-1 process creating double layers of Ge (a) and of the standard MIC-2 process, known as ALILE[Nast, 2000b].

Figure 2 .

 2 Figure 2.37 summarizes the structures generated by the newly identified MIC-1 mechanism (Figure 2.37 (a)) and by the standard MIC-2/ALILE process (Figure 2.37 (b)).

Figure 2 .

 2 Figure 2.38 -Ex situ Nomarski optical images of samples at different stages of annealing. A,B: sample before and after the appearance of MIC-1 dendrites. C: Fully annealed sample. Scale bars: 10 microns.

Figure 2 .

 2 Figure 2.39 -GIXRD of the samples shown in Figure 2.38. The top axis indicates the position of Ge-related diffracted intensities. Red labels mark the GIXRD peaks characteristic of a perfect (111) texture. 220, 311, 222) are indicated by arrows. A 0 is the diffraction of a single al layer. Letter R refers to a reference Ge powder diffractogram. The baseline of each diffractogram is indicated by a solid black line.

  GIXRD . MIC-1 sample B shows a weak preferential (100) texture while the fully annealed sample C is more textured along the (111) direction.

Figure 2 .Figure 2 .

 22 Figure 2.40 -(a) Schematic of the working configuration for an EBSD detector in a SEM; (b) Example of diffraction pattern with Kikuchi bands of Ge; (c) Schematics of diamond crystal structure with three main crystalline directions and their color code in our EBSD maps.

  Ge layer obtained by MIC-2/ALILEDescription of the experiment Using EBSD, we investigate the local crystal orientation of each grain over a wide area of the sample characterized by GIXRD in the previous section 2.6.1.

Figure 2 .

 2 Figure 2.41 -EBSD analysis of the Ge layer after complete AIC and removal of Al by wet etching (sample C of Table 2.6). The normal of the Ge crystallites is shown projected along the vertical direction (i.e. perpendicular to the substrate) in (a), and in-plane(parallel to the substrate) in (b).

Figure 2 .

 2 Figure 2.42 -EBSD analysis of the sample (28F9) partially annealed with coexisting MIC-2/ALILE and unreacted areas. The sample is not etched. IQ image of the MIC-2/ALILE area on the left and unreacted area on the right (a). The normal of the Ge crystallites, for the same area, is shown projected along the vertical direction in (b). EDXS maps of Al and Ge distribution are shown in (c) and (d), respectively. The unreacted a-Ge is filtered out from the images and thus appears back.

Figure 2 .

 2 Figure 2.43 -EBSD analysis of the Ge layer after complete AIC and removal of Al by wet etching. The normal of the Ge crystallites is shown projected along the vertical direction in (a) and in-plane in (b)).

Figure 2 .

 2 Figure 2.44 -(a)-(e) Schematics of the new nucleation by MIC-2/ALILE (black arrows) which occurs as the crystallization front touches an existing Ge dendrite. The different colors indicate the different in-plane orientations (see reference in-plane EBSD map in 2.41). (f)-(l): Selected views of in situ optical monitoring of a MIC sample from the substrate side. The possible nucleation locations are indicated by red arrows. Annealing times for are 6.07 h, 6.43 h, 6.80 h, 7.12 h and 7.25 h from (f) to (l). Scale bar is 20 µm for (f)-(l). (k) EBSD map of in-plane direction of typical MIC sample with different in-plane orientations around the central dendrite.

  Photoresisit

Figure 2 .

 2 Figure 2.46 -Schematic of the sample preparation process: spin of resist (AZ5214) and 1 min soft bake at 125 °C (a); exposure with an inverted mask for 5 s, 1 min reversal bake at 125 °C, flood exposure 30 s without the mask, developing (developer AZ826) (b); e-beam deposition of the stack Ge/Al/AlGeOx/Ge (c); lift-off in acetone (d); low temperature 250-300 °C annealing (e); selective Al wet etching of the top Al layer (f).

Figure 2 .Figure 2 .

 22 Figure 2.47 -(a) Nomarski images of patches with four different side lengths (40, 90, 175 and 275 µm), after full annealing and processing. (b): SEM image of a square patch of 40 µm side; the MIC-1 dendrites are indicated by black arrows.

Figure 2 .

 2 Figure 2.49 -Total area density of the Ge dendrites as a function of the patch area. Insert shows how the total density is computed from the edge and area dendrite densities. The blue dashed line is the the total density d tot , computed from equation 2.1 as function of the area A.

Figure 2 .

 2 Figure 2.50 -In situ optical images of square patterns (with 2.7 µm wide patches) during the annealing, before (a) and after (b) the appearance of the Ge dendrites due to the MIC-1 process. Red arrows indicate some of the Ge dendrites.

Figure 2 .

 2 Figure 2.51 -In situ optical images of the top of the sample (front-side, sample 38O1) during the AIC annealing at 248 °C. (i) and (j) are zoomed views of the areas indicated by rectangles in (c) and (d), respectively. Scale bar for images (a) to (h) is 40 µm. Black and gray arrows indicate the MIC-1 Ge dendrites on the edge and in the center of the pattern respectively. Red arrows indicate the onset of the inversion process (MIC-2/ALILE) when Al moves to the top layer.

Figure 2 .

 2 Figure 2.52 -(a) In-situ optical image obtained during annealing of patterned sample. (b) Corresponding particle identification and counting obtained by ImageJ after the appropriate thresholding of the initial image.

Figure 2 .

 2 Figure 2.53 shows the measured number of crystallized patches as a function of the annealing time. The graph in Figure 2.53 is normalized over the ∼1700 patterns observed. Inserts show in-situ optical images at specific durations of the annealing.

Figure 2 .

 2 Figure 2.53 -Normalized count of the 2.7 µm patches crystallized via MIC-2/ALILE, as a function of time. Left axis: annealing temperature. Inserts show selected optical images at three specific times (6 h, 19.5 h, 29.5 h). The red line is an exponential fit of the fraction of crystallized patterns, while the green line is a sigmoidal fit.

Figure 2 .

 2 Figure 2.54 -(a)-(e) In situ optical images of a patterned sample (38O1, already shown in Figure 1.51) during annealing.(f) Nomarski optical image after Al wet etching. Comparison of EBSD analysis of the Ge patch (g,i) and a reference non-patterned sample (see section 2.6.2.1 p. 57) (h-l). EBSD in the vertical direction (g-i) and in-plane direction(h-l)), i.e. parallel to the substrate.

Figure 2 .

 2 Figure 2.55 -SEM image (a) and EBSD maps of the out-of-plane (b) and in-plane (c) crystal orientation of the final Ge layer in a 2.7 microns square patch.

Figure 2 .

 2 Figure 2.56 -Nomarski images of the top Ge layer at the bilayer edge (a), top Ge layer edge (b) and Al layer edge(c), as schematized at the bottom.

Figure 2 .Figure 2 . 59 -

 2259 Figure 2.58 -In situ optical image of the top of the sample (front-side, sample 84) during the AIC annealing at 225-260 °C (two temperature annealing). Scale bar for images (a) to (h) is 40 µm. White arrows indicate the MIC-1 nucleation and growth; purple arrows indicate the MIC-1 nucleation outside the pattern and red arrow the nucleation of the MIC-2/ALILE

  Figures 2.60

  (a) and (b) show Nomarski optical images of a patterned sample obtained using the revised lithography process (a) and of a reference sample with a continuous layer (b), after the annealing and the wet etching of the top Al layer. The deposition sequence, identical for the two samples, corresponds to the previously optimized structure (Ge/Al/GeOx/Ge, 1/20/1.5/20 nm, with 12 min air oxidation of the interlayer). The two samples are fabricated in the same batch of e-beam deposition and annealed side by side at 245 and 265 °C. The pattern is an array of squares of different sizes (used in the previous study, seeFigure 2.47 p. 67). We observe comparable features in both the patterned and continuous samples, with the growth of large and flat Ge platelets. The patterned sample exhibits a lower coverage in the patterned areas compared to the reference continuous layer.

Figure 2 .

 2 Figure 2.60 -(a,b) Ex situ optical images of Ge layers obtained after Al removal. Top Ge layer patterned sample (a) and continuous layers sample (b). (c) GIXRD diffractograms of samples shown in (a,b), compared to reference powder Ge data (randomly oriented). Reflection indices are indicated at the top; the 220 422 and 440 (red labels) are characteristic of the (111) texture.

Figure 3 . 1 -

 31 Figure 3.1 -Diamond cubic crystal structure unit cell of Ge. Lattice parameter is a=5.658 Å.

Figure 3 . 2 -

 32 Figure 3.2 -Unit cell of the diamond cubic crystal of Ge projected in different orientations.

Figure 3 . 3 -

 33 Figure 3.3 -Picture of the Si-G MBE reactor.The chamber is composed by: the main viewport, the viewport to see the Si crucible, the Medium High Temperature (MHT) Ge cell, the manipulator (external and internal view), the gauge pressure sensor, the RHEED viewport with phosphorous screen, the electron gun for Si evaporation, the pyrometer, the quartz sensor, the quartz oscillator, and a free cell port. On the manipulator we see: the rotating parts, the oven and the molybdenum sample holder.

Figure 3 . 4 -

 34 Figure 3.4 -(a-b): cross-section SEM characterization of the sample for the Ge deposition rate calibration at two different zones. (c): corresponding chamber pressure as a function of the Ge cell temperature. (d): measured SEM thickness in (a,b), converted into growth rate and quartz rate as function of the pressure inside the chamber.

Figure 3 .

 3 Figure 3.5 -(a) Example of RHEED oscillation of a growing Ge layer on a Ge(100) wafer using a Ge cell temperature of 1300 °C; (b) RHEED image with red square corresponding to the analyzed area for the oscillation. The substrate temperature is 300 °C.

Figure 3 . 6 -

 36 Figure 3.6 -Picture of the equipment used for thermal deoxidation and XPS characterization.

Figure 3 . 7 -

 37 Figure 3.7 -XPS characterization of MIC-Ge sample at each annealing step showing the Ge 3d and Ge 2p peaks respectively at ∼30 eV (a) and ∼1218 eV (b); the spectra of the O 1s peak (c); the spectra of the C 1s peak (d).

Figure 3 . 8 -

 38 Figure 3.8 -Schematics of MIC fabrication and the epitaxy of Ge on the Ge platelets crystallized on SiO 2

Figure 3 . 9 -

 39 Figure 3.9 -Ge desorption flux (F d ) in monolayers per second (ML/s) as a function on the substrate temperature. Black circles and black line are the threshold values for the formation of stable Ge nuclei on thermal SiO 2 . Figure extracted from ref.[Li, 2006b]

Figure 3 .

 3 Figure 3.10 -SEM characterization of Ge grown by MBE on MIC-Ge pseudo-substrate on silica with low Ge flux (Ge cell at 1200 °C) for ∼2 h 20 min. Top view image (a) and 45°view (b,c).

Figure 3 .

 3 Figure 3.11 -SEM characterization of Ge epitaxy on MIC-Ge sample: with Ge cell at 1200 °C for ∼2 h 20 min (a,c) and with Ge cell at 1240 °C for ∼1 h (b,d). Top view image (a,b) and 45°view (c,d).

Figure 3 .

 3 Figure 3.12 -EBSD analysis of the epitaxial growth Ge layer on MIC-Ge layer (high Ge flux (56B) of Figure 3.11). The normal of the Ge crystallites is shown projected along the vertical direction (a), i.e. perpendicular to the substrate, and in the plane (b), i.e. parallel to the substrate.

Figure 3 . 13 -

 313 Figure 3.13 -Unit cell of the GaAs cubic zinc blende crystal structure. Lattice parameter is a 0 = 5.658 Å.

Figure 3 .

 3 Figure 3.14 -Atomic stacking of GaAs in different crystal directions: (a) <001>, (b) <111>B, (c) <111>A

Figure 3 .

 3 Figure 3.15 -The III-V MBE reactor Riber32 at C2N.

Figure 3 .Figure 3 . 16 -

 3316 Figure 3.16 shows the schematics of the Riber32 MBE chamber. On one side are the effusion cells pointing to the sample at the center of the chamber. Each cell aperture is fitted with a shutter controlling the flux toward the sample. A RHEED gun directs an electron beam toward the sample with a grazing angle. The electron beam impinges a phosphorous screen on the opposite side of the chamber after it has interacted with the sample surface. A pyrometer points to the sample surface measures its temperature during the growth. The sample is introduced in the chamber using a mobile arm. Several pumps keep the pressure low in the growth

Figure 3 . 17 -

 317 Figure 3.17 -Measured Ga flux (BEP) as a function of the inverse of the Ga cell temperature in log scale. The red line is linear fit of the experimental values (squares).

Figure 3 .

 3 Figure 3.19 -GaAs (100) growth rate at 580 °C as a function of the measured Ga flux (BEP) from Equation 3.2

Figure 3 .

 3 Figure 3.22 -GaAs (100) growth rate as a function of the As flux (BEP) in the As-limited regime.

Figure 3 .

 3 Figure 3.23 -RHEED reconstruction of the GaAs(111)B surface observed: before growth at 520 °C (a), after growth at 520 °C (b) and after growth at 400 °C (c). Figure extracted from[Cho, 1970]

Figure 3 .

 3 Figure 3.24 -(a) Map of the RHEED reconstructions at different substrate temperatures and arsenic fluxes from GaAs(111)B (from[Woolf, 1993a; Woolf, 1993b]). (b) Table of RHEED reconstructions as a function of the growth temperature (from[Woolf, 1993a]).

Figure 3 .

 3 Figure 3.25 -(a) Map of the RHEED reconstructions at different substrate temperatures and arsenic fluxes from GaAs(111)A (from[Woolf, 1993b]). (b) Table of RHEED reconstructions as a function of the growth temperature (from[Woolf, 1993a]).

Figure 3 .

 3 Figure 3.26 -(a): Surface morphology of 1 µm GaAs on GaAs(111)B substrate growth at 550 °C, 610 °C and 660 °C (from[Tsutsui, 1990]). (b): Nomarski characterization of GaAs and AlGaAs structure grown on GaAs(111)B substrates as function of the As flux.( from[Yeo, 2000]).

Figure 3 .

 3 Figure 3.27 -Growth series of GaAs(111)A layer grown using different homoepitaxy in As 4 and Ga flux (a) and different in growth temperature (b). Surface morphology defects (P: pyramidal; I: irregular) observed as a function of flux ratio and substrate temperature (c). SEM characterization extracted from Ritzmann et al.[Ritzmann, 2018]. Summary of the morphology and defect from and resuming graph from Sato et al.[Sato, 1994].

Figure 3 .

 3 Figure 3.28 -RHEED pattern of GaAs layer grown on a GaAs (111)B wafer: (a) before growth at T=637 °C with a 1×1 pattern; (b) after growth, with substrate temperature of 520 °C with a 2×2 reconstruction; (c) after growth, with a √ 19 × √ 19 reconstruction (As 4 BEP: •10 -6 Torr; Ga BEP: 1.44•10 -7 Torr)

Figure 3 .

 3 Figure 3.29 -AFM characterization (top) and Nomarski optical characterization (bottom) of GaAs(111)B layer grown using As 4 :Ga BEP ratios from 11:1 to 51:1 and a substrate temperature of ∼620 °C. The AFM images are 5×5 µm wide. Nomarski image are taken using three different magnitudes: 100×, 50× and 20×.

Figure 3 .

 3 Figure 3.30 -AFM characterization (left) and Nomasrki optical characterization (right) of GaAs(111)B layer grown using substrate temperature of ∼580 °C, ∼600 °C, ∼620 °C, ∼630 °C and ∼640 °C. As 4 :Ga BEP ratio of 42-49:1. The AFM image are either 5×5µm or 20×20µm. Nomarski image are taken at three different magnitudes: 100×, 50× and 20×.

Figure 3 .

 3 Figure 3.31 -Surface morphology of GaAs(111)B layer grown using substrate temperature of 580-640 °C, As 4 :Ge BEP ratio from 11:1 to 90:1 and (Ga BEP of 1.4 -1.5 • 10 -7 Torr). The blue rectangle indicates the As 4 series described in Figure 3.29. The red rectangle indicates the temperature series characterized in Figure 3.30. The dashed yellow rectangle highlights the smoothest layer.

Figure 3 .

 3 Figure 3.32 -Photoluminescence spectra of GaAs grown on GaAs(111)B substrates at 580 °C, 600 °C, 620 °C and 640 °C.

Figure 3 .

 3 Figure 3.33 -Nomarski optical images of GaAs homoepitaxy on GaAs(111)B made with the same As valve aperture: As 4 (samples #5 and #8) or As 2 (samples #13 and #14), at 600 °C or 640 °C.

Figure 3 .

 3 Figure 3.33 compares Nomarki images of GaAs grown with As 4 and As 2 molecules at 600 °C and 640 °C.

Figure 3 .

 3 Figure 3.34 -Schematics of an antiphase domain (APD) in GaAs grown on Ge(100).

Figure 3 .

 3 Figure 3.34 shows how APDs can form due to growth starting on adjacent terraces separated by a single atomic surface step on the Ge (001) surface. The two phases

Figure 3 . 35 -

 335 Figure 3.35 -Schematic of GaAs growth on Ge(111) with a full step.

Figure 3 .

 3 Figure 3.36 -(a): schematic of the Ge(111) surface terminated As with lone pair electrons. (b) Density of state bulk non-reconstructed surface and an As passivated surface, Figure extracted from[Bringans, 1992]

Figure 3 .

 3 Figure 3.37 -Schematic of the GaAs inverting polarity from (111)A to (111)B by the insertion of a Ge interlayer. Figure extracted from[Koh, 2001]

Figure 3 . 38 -

 338 Figure 3.38 -Schematics of the formation of a twin at the GaAs/Ge interface

Figure 3 .

 3 Figure 3.39 -RHEED characterization of GaAs layer on Ge(111) wafer: (a) before growth; (b) after growth, with substrate temperature of 550 °C; (c) after growth, with substrate temperature of 600 °C. Observed reconstructions: 1×1 (a), 2×2 (b) and (c)

Figure 3 .

 3 Figure 3.40 -AFM characterization ∼10 nm (a) (sample 77114) and ∼450 nm (b) (sample 76745) of GaAs growth on Ge at 600-610 °C. (c) AFM zoom of a pyramid of the thick sample. (d) profile of the pyramid AFM image shown in (c).

Figure 3 . 41 -

 341 Figure 3.41 -GaAs layer grown on Ge(111). Reciprocal space maps around the 111 (a) and 331 (b) reflections. The red dotted line marks the peak alignment along q z .

Figure 3 .

 3 Figure 3.42 -High Resolution X-ray Diffraction in triple axes configuration of the 111 reflection. The two peaks correspond to the Ge and the GaAs peaks.

Figure 3 .

 3 43 shows Nomarski images of samples grown at different temperatures T and under different As fluxes. The main results are the following: • sample A (low T, low As flux) displays high pyramids; • sample B (low T, intermediate As flux) has a relatively smooth surface with isolated defects;• sample C (low T, high As flux) is rougher, but without pyramids;• sample D (high T, low As flux) displays pyramids;• samples E and F (high T, intermediate and high As fluxes) have relatively smooth surfaces;

Figure 3 . 43 -

 343 Figure 3.43 -Nomarski characterization of GaAs on Ge(111) 450 nm grown with an As:Ga BEP ratio of: 37:1 (a,d), 63:1 (b,e) and 94:1 (c,f). Sample grown at 600-610 °C (a-c) and at 630-640 °C (d-f).

Figure 3 .

 3 Figure 3.44 -SEM top view and tilted views of GaAs on Ge(111) 450 nm grown with an As:Ga BEP ratio of: 37:1 (a,d), 63:1 (b,e) and 94:1 (c,f). Sample grown at 600-610 °C (a-c) and at 630-640 °C (d-f).

Figure 3 . 45 -

 345 Figure 3.45 -Schematic of the rotated pyramids: some pyramidal hillocks are rotated by 180°with respect to the dominant direction of the pyramids.

  2 and 2 s for Ga) or by the type of As molecule used (As 4 at 7•10 -6 Torr BEP or As 2 at 5•10 -6 Torr). The substrate temperature is kept constant at 600-610 °C. SEM cross-section views of the ∼45 nm thick GaAs layer grown on Ge.

Figure 3 .

 3 Figure 3.47 (b,e) shows tilted SEM views (45°) of the same sample (As 4 preexposure). The pyramids are uniformly distributed all over the surface. For a similar sample grown with As 2 (Figure 3.47 (c,f)), the density of pyramids is slightly lower. If growth is started with Ga pre-exposure, Figure 3.47 (a,b), we observe a large density of pyramids rotated by 180°.

Figure 3 .

 3 Figure 3.47 -SEM characterization at 45°of ∼45 nm GaAs on Ge(111) growth. Growth started with Ga (a) and with As 4 (b). Growth started and continued with As 2 (c).

Figure 3 .

 3 48 shows 45°-tilted SEM images of the GaAs layers grown without (a,d) and with (b,e) type I MEE layer. With MEE, the pyramids are still present but a little flatter.

Figure

  Figure 3.48 (c) shows the sample with the type-II MEE layer. We observe the appearance of additional small pyramids but the morphology is similar.

Figure 3 .

 3 Figure 3.48 -SEM characterization at 45°of ∼45 nm (5 min) GaAs on Ge(111) growth starting with Ga (a,d). SEM of the Ga buffer layer as in (a) with an additional MEE layer of As 4 /Ga 4/4 s for 5 min (b,e) and Ga/As 4 /close 2/4/2 s for 76 times (c-f). For each sample two magnitude are shown.

Figure 3 .

 3 Figure 3.49 -SEM characterization in top view, cross-section view and tilted view of GaAs growth on Ge(111) with thickness of ∼45 nm, ∼450 nm and ∼1.65 µm.

  experiment We performed X-ray pole figure characterization using the 111 pole of three samples grown with different As fluxes (BEPs of 4.3•10 -6 , 8.9•10 -6 and 1.35•10 -5 Torr) at a substrate temperature of 600-610 °C.ResultsFigure 3.50 shows the reference 111 pole figure of the Ge substrate and that of the three samples made at low, intermediate and high As 4 flux (see section 3.4.2.4) combined with top view optical images. • The pole figure of the reference Ge (111) substrate shows a symmetry of order 3, as expected for a single crystal. • At intermediate As flux (Figure 3.50 (b)), we observe the main 3 peaks and three other peaks possibly resulting from on-axis (growth plane) twinning (rotation R1) in the GaAs layer. • The sample made with low As flux (Figure 3.50 (c)) shows additional reflections, possibly from other simple twinning on the (111) planes different from the growth plane (rotations R2, R3, R4). • At high As flux, peaks indicative of multiple twinning are clearly visible (rotation combinations R1R4, R1R2, R1R3).

Figure 3 . 50 -

 350 Figure 3.50 -Pole figure around of the 111 reflections of: (a) the reference Ge substrate, (b) GaAs made with intermediate As 4 (BEP of 8.9 • 10 -6 Torr), (c) lower As 4 (BEP of 4.3 • 10 -6 Torr), (d) higher As 4 (BEP of 1.35 • 10 -5 Torr). Inserts are the Nomarski characterization of the surface morphology.

Figure 3 . 52 -

 352 Figure 3.52 -Optical view (Nomarski) of the small defective area (circle white dashed) selected for the cross-sectional TEM analysis (a). The yellow line in (a) indicates the position of the lamellae extracted with the FIB. (b) schematic of the sample structure.

Figure 3 .

 3 Figure 3.53 -Dark field TEM cross-section (a) inside the defective area and (b) outside. (c) HRSTEM-HAADF image of the boundary of the twinned domain. (d) Schematic view of the twin boundary.

Figure 3 .

 3 Figure 3.54 -SEM images (a,c) with corresponding EBSD maps (b,d)). The area imaged in (c,d) is indicated by a red square in (a,b) (e) EBSD Euler angle map.

Figure 3 .

 3 Figure 3.55 -SEM (a,c) and corresponding CL (b,d) images of an area with a twinned domain (a,b). Red arrows in (a,b) indicate the inverted pyramids and so the twinned area. In (c), red and blue indicate the possible twin orientation (alpha and beta). Purple lines in (d) indicate where the CL signal is very low. (e) Cross-section schematics of the incoherent interface (purple dashed line) and the coherent interface in the (111) plane (blue dashed line) created by a twin.

Figure 3 .

 3 Figure 3.56 (d) shows an intensity profile along the line marked in yellow in. On average, the dumbells are slightly more intense then on the left,

Figure 3 . 56 -

 356 Figure 3.56 -Schematic of a GaAs(111)B (a) and GaAs(111)A (b) polarity, from side view. HAADF image of the GaAs on Ge (c). (d): intensity profile of the yellow line in (c).

Figure 3 .

 3 Figure 3.57 -STEM-EDX of GaAs on Ge(111) cross-section view (a); relative HAADF image (c) and scheme of crystal structure corresponding to the EDX map (b). Schematic of the cross-section view of the GaAs with respect to the Ge substrate (d).

Figure 3 .

 3 58 (b,c) show the chemical composition profile across the layers. The zero position is set at the initial Ge/GaAs interface. (Figure3.58 (c) is a zoom on Figure3.58 (b)).

Figure 3 .

 3 Figure 3.58 -(a) HAADF of cross-section view showing the Ge(111) substrate, the main GaAs layer and the AlGaAs marker layer. (b,c): corresponding EDX chemical composition along the white line in (a) with two different scale, EDX element map of Ge (d), As (e), Ga (f), Al (g).

Figure 3 .

 3 Figure 3.59 -SEM 45°view of MIC-Ge islands on the silica substrate: after etch of the Al (a,b) and after additional cleaning with H 2 O 2 and HF (c,d).

Figure 3 .

 3 Figure 3.60 -SEM characterization of 45°tilted samples of GaAs growth on MIC-Ge substrate at 600 °C, 620 °C, 635 °C, 645 °C and 670 °C.

Figure 3 .

 3 Figure 3.61 -SEM characterization of GaAs on MIC-Ge with different thickness from left to right: only Ge; GaAs ∼45 nm; GaAs ∼360 nm; GaAs ∼720 nm; GaAs 2 µm;. For each sample: tilt view, top view (45 °) and cross-section view.

Figure 3 .

 3 Figure 3.62 -SEM tilt view (45°) at two different magnitude of GaAs on MIC-Ge with different thickness made with As 2 .

Figure 3 .

 3 Figure 3.63 -Large scale SEM-SE (a,d) characterization of MIC-Ge layer (a-c) and epitaxial GaAs (d-f). EBSD characterization in the vertical direction (b,e) and in the plane (c,f).

Figure 3 .

 3 Figure 3.64 -Small scale characterization of a MIC-Ge layer (a-c) and epitaxial GaAs (d-f). (a,d) SEM; (b,e) EBSD crystal orientation in the vertical direction; (c,f) in-plane.

Figure 4 . 1 -

 41 Figure 4.1 -Schematic of the identical solar cell structures to be grown on three different substrates: homoepitaxy on GaAs (111) A and B wafers, heteroepitaxy on commercial Ge (111) wafers and MIC-Ge pseudo-substrates.

Figure 4 . 2 -

 42 Figure 4.2 -Picture of the Riber Compact 21 (a) and the Riber 32 (b) MBE machines used to grow solar cells structures.

Figures 4. 3

 3 Figures 4.3(a) and (b) respectively show the details of the GaAs-based solar cells grown by Wang et al.[Wang, 2017] and Venkatasubramania et al.[Venkatasubramanian, 1996]. Both cells have a similar structure, with a buffer/contact layer, a back-surface field (BSF) layer, a thick n-doped GaAs base, a thin p-doped GaAs emitter, a window layer and a final p+ top contact layer.

  Figures 4.3(a) and (b) respectively show the details of the GaAs-based solar cells grown by Wang et al.[Wang, 2017] and Venkatasubramania et al.[Venkatasubramanian, 1996]. Both cells have a similar structure, with a buffer/contact layer, a back-surface field (BSF) layer, a thick n-doped GaAs base, a thin p-doped GaAs emitter, a window layer and a final p+ top contact layer.

Figure 4

 4 Figure 4.4 -N-type carrier densities of GaAs layers doped with silicon impurities for the GaAs(111)B (black full circle) and GaAs(100) (empty circles) as a function of the temperature of the Si effusion cell (from[Woolf, 1992]).

Figure 4 . 5 -

 45 Figure 4.5 -Carrier concentration and doping type in Si doped GaAs(111)A layer grown at various substrate temperatures, as a function of the As 4 :Ga flux ratio (from[Sato, 1996]). Degree of compensation of the Si impurities in GaAs layers, (111)A (circle), (111)B (triangle) and (100) (cross) (b), extracted from[Yaremenko, 2008].

Figure 4 . 6 -

 46 Figure 4.6 -Schematic of the test structure used for the carrier density measurement.

Figure 4

 4 Figure 4.7 -N-type carrier density (cm -3 ) as a function of the temperature of the Si cell in GaAs(111)B compared to reference GaAs(100). Substrate temperature is ∼600 °C (black squares) and ∼640 °C (green circle and red triangle). Square and circle samples are grown using As 4 , while the triangle sample is grown using As 2 .

Figure 4

 4 Figure 4.8 -P-type carrier density (cm -3 ) as a function of the temperature of the Be cell in GaAs(111)B, compared to reference GaAs(100). Substrate temperature of ∼600 °C. We use As 4 only.

Figure 4 . 9 -

 49 Figure 4.9 -Nomarski characterization at three different optical magnitude (100x, 50x and 20x) of an undoped GaAs(111)B layers (left), Si-doped sample (center) and Be-doped (right).

Figure 4 . 10 -

 410 Figure 4.10 -Full structure of the solar cell on a GaAs wafer, showing the bottom and top contacts, as well as the details and doping type of each layer.

Figure 4 .

 4 Figure 4.11 (c) shows a cross-sectional viewed by SEM. The bottom and top AlGaAs layer are visible as thin dark lines. We measure a total thickness from the bottom AlGaAs BSF to the top surface of ∼1300 µm, instead of ∼1420 µm expected. Inside the middle GaAs layers, we see a different SE contrast from the dark p-type doped layers (top) to the light n-type doped layers (bottom).SEM top viewSEM cross-sectional

Figure 4 .

 4 Figure 4.11 -SEM 45°view (a) and Nomarski optical characteristic at 100x magnitude (b) of the GaAs solar cell surface, solar cell sample #1 (136). Cross-sectional SEM view (c) of the structure.

Figure 4 . 15 -

 415 Figure 4.15 -IV curve for the closest contacts before (a) and after (b) annealing (at 380-400 °C 1 min) of the Ti/Au top contact on of the top contact GaAs p-type before and after annealing. Measured resistance for the different distances between the TLM patterns before and after annealing (c).

  Figures 4.16

  (a) and (b) show typical I-V curves before and after annealing. We observe a very resistive behavior with non-ohmic component before annealing and a pure ohmic behavior of the contact after the annealing.

  Figure 4.16 -Figures 1.16(a) and (b) show typical I-V curve of the contact on the backside of n-type Ge (111) wafer before and after annealing 1 min at 380-400 °C. Measured resistance for the different distances between the TLM patterns after annealing (c).

  Figure 4.16 -Figures 1.16(a) and (b) show typical I-V curve of the contact on the backside of n-type Ge (111) wafer before and after annealing 1 min at 380-400 °C. Measured resistance for the different distances between the TLM patterns after annealing (c).

Figure 4 . 17 -

 417 Figure 4.17 -IV curve for the closest contacts before annealing (a) and after annealing (b) for the back contact Ge n-type. Measured resistance for the different distances between the TLM patterns (c).

Figure 4 . 18 -

 418 Figure 4.18 -IV curve (current density) of a GaAs solar cell of the sample #1 (136) in light and dark condition. (b). Calculated power density from the IV curve in the light condition

Figure 4 . 19 -

 419 Figure 4.19 -External quantum efficiency of a GaAs solar cell on GaAs(111)B (sample 136).

Figure 4

 4 Figure 4.20 -SEM bird view 45°(a), cross-section (b) and Nomarski optical microscopy (c) of the solar cell.

Figure 4 .

 4 Figure 4.21 -IV curve of a GaAs solar cell of the cell sample #2 (140) in light and dark condition with comparison of the IV curve shown in section 4.4.3.1, p. 170.

Figure

Figure 4 .

 4 Figure 4.22 -SEM 45°(a) and Nomarski optical characteristic at 100× magnitude (b) of the solar cell sample #3 (143). SEM cross-sectional view (c).

Figure 4

 4 Figure 4.23 -IV curve of a GaAs best solar cell of sample #3 (143) in light and dark condition with comparison of the IV curve of the best solar cell of section 4.4.3.1, p. 170

Figure 4 .

 4 Figure 4.25 -SEM (45°tilt) and optical (top view) images at high and low magnification of the top surface of full solar cell structure on different substrates: (a) reference GaAs(111)B, (b) GaAs(111)A, (c-d) Ge(111) grown at T sub =620 °C or T sub =640 °C.

Figure 4 . 27 -

 427 Figure 4.27 -Nomarski optical microscopy of the MIC-Ge substrate before GaAs epitaxy.

Figure 4 .

 4 Figure 4.28 -(a-b) SEM (45°tilt) and (c) optical (top view) images of the top surface of full solar cell structure on the MIC-Ge pseudo substrate. (d) SEM cross-section of the obtained layers.

  Figure A.1 -Scheme of the XRD θ -2θ configuration (cross sectional view).

  Figure A.2 (a)). For a full (111) crystal texture, the [111] direction is normal to the surface, while the in-plane orientation is random (Figure A.2 (b)).

  Figure A.3 (a) shows relative XRD intensities for a Ge powder and Figure A.3 (b) for a Ge (111) texture, as a function of 2θ.

Figure A. 2 -

 2 Figure A.2 -Schematic of the possible crystal orientations in a powder structure (a) and in a fully textured powder sample (b).

Figure A. 3 -

 3 Figure A.3 -Relative XRD intensities for a (random oriented) Ge powder (a) and a fully (111) textured Ge powder (b) in the standard (θ-2θ) configuration.

  Figure A.4 -Cross view (a), top view (b) and 3D view (c) of the GIXRD configuration. Intensities are collected as a function of the in-plane angle (2θχ) between the plane of the incidence bean and the detector position in the plane. The plane of incidence beam is marked as a red rectangle in (c).

Figure A. 5 -A. 3 . 2

 532 Figure A.5 -Picture of the RIGAKU Smartlab diffractometer at C2N.

I

  220 + I 422 + I 440 all I ijk

Figure A. 6 -

 6 Figure A.6 -Ge crystal unit cell with vertical [111] direction.

Figure A. 7 -

 7 Figure A.7 -Hexagonal symmetry in the (111) plane.

Figure A. 8 -

 8 Figure A.8 -Four-fold symmetry in the (100) plane.

  GIXRD for the three different samples. The values are reported in Figure A.4. For a perfect (111) texture, the L (100) GIXRD factor has high value of 0.65. Similarly, for a perfect (100) texture the L (111) GIXRD factor reaches 0.83. This is not the case in standard XRD where computing the L (100) GIXRD for a pure (111) textured sample would result in a negative values.Therefore, the constructed Lotgering factor for GIXRD is less sensitive than the standard XRD Lotgering factor due to overlapping reflection families. As seen above the diffracted intensities I 220 and I 440 contributes to both L the discriminative peaks have a very low intensity compared to the 220 reflection.

Figure A. 9 -

 9 Figure A.9 -Experimental X-ray spectrogram of MIC-Ge sample (see chapter 2, section 2.6.1) in GIXRD configuration, after Al etching.

  GIXRD of the experimental sample for the (111) and (100) texture.

  Figure B.1 -Résumé graphique de la thèse : fabrication de pseudo-substrats MIC-Ge sur silice, épitaxie de GaAs sur MIC-Ge et finalement cellules solaires complètes.

Figure B. 2 -

 2 Figure B.2 -Photographie et schémas de la plaque chauffante équipant le microscope optique in situ.

Figure B. 3 -

 3 Figure B.3 -Microscopie optique in situ. Vue de l'arrière de l'échantillon lors du recuit AIC-Ge à 238 °C. (q) et (p) sont des vues agrandies des zones indiquées par des rectangles en (d) et (i), respectivement. La barre d'échelle est 40 µm. Les flèches noires indiquent la nucléation et la croissance de MIC-1. Les flèches rouges indiquent le début du processus MIC-2. (o) Schéma du MIC-1 et du MIC-2.

  Figure B.4 (bleu)), chacun organisé autour d'une zone dendritique dont la normale tend à être selon (100) (Figure B.4 (rouge)).

  Figure B.4 -Analyse EBSD de la couche de Ge après AIC complète et attaque de l'aluminium par gravure humide. La direction normale des cristallites de Ge est représentée projetée le long de la direction verticale (c'est-à-dire perpendiculaire au substrat) en (a) et dans le plan (parallèle au substrat) en (b).

Figure B. 5 -

 5 Figure B.5 -Caractérisation en microscopie optique de couches de GaAs épitaxiées sur Ge (111) (épais d'environ 450 nm) avec des rapports de flux équivalents (BEP) As :Ga de 37 :1 (a,d), 63 :1 (b, e) et 94 :1 (c,f). Croissances à 600-610°C (a-c) et à 630-640°C (d-f).

  Nous caractérisons ensuite l'orientation cristalline de la couche de GaAs par EBSD. La Figure B.6 montre des images de la couche MIC-Ge après l'épitaxie de GaAs. Les couches de GaAs conservent la texture cristalline initiale (111) et la géométrie dans le plan caractéristique des couche de MIC-Ge, ce qui démontre le bonne épitaxie des couches de GaAs sur les îlots de Ge.

Figure B. 6 -

 6 Figure B.6 -Caractérisation de la couche de GaAs épitaxiée sur les pseudo-substrats MIC-Ge (épais d'énviron 45nm) : images de microscopie à balayage en électrons secondaires (a), EBSD dans la direction verticale (b) et dans le plan (c).

B. 4

 4 Cellules solaires sur GaAs (111) A et B et MIC-GeNotre objectif final est de fabriquer des cellules solaires complètes basées sur un absorbeur de GaAs, en utilisant les pseudo-substrats MIC-Ge précédemment étudiés. A cet effet, nous effectuons tout d'abord la croissance de cellules solaires en GaAs sur différents substrats monocristallins : GaAs(111)B, GaAs(111)A et Ge(111). Sur GaAs (111)B, nous définissons les paramètres de la cellule solaire, tels que les dopages p et n et déterminons les conditions de croissance requises. Les cellules sont caractérisées par leur réponse courant-tension (IV) et leur réponse spectrale (EQE). Pour obtenir une meilleure efficacité, nous déposons un revêtement antireflet (ARC) sur la meilleure cellule et obtenons un rendement d'environ 16% . Les caractérisations sont montrées dans la Figure B.7.

  Figure B.7 -(a) Caractérisation I(V) sous éclairement et dans l'obscurité, avec et sans couche antireflet (ARC), de notre meilleure cellule solaire. (b) EQE et réflectivité avec et sans ARC.

Figure B. 8 -

 8 Figure B.8 -Images en microscopie électronique à balayage (inclinaison à 45 °) de la surface des cellules solaires fabriquées sur différents substrats : (a) GaAs(111)B (référence), (b) GaAs(111)A ; (c,d) Ge(111) à T sub = 620 °C ou T sub = 640 °C.

  ) et(111). L'étape MIC-2 est le processus standard d'inversion des couches et produit majoritairement des cristaux de Ge avec une texture (111) ; voir Figure (B.9).

  Figure B.9 -Schémas du processus MIC-1, qui crée une couche de Ge cristallin (c-Ge) en dessous de la couche initiale de Ge amorphe (a-Ge)[START_REF] Pelati | In Situ Optical Monitoring of New Pathways in the Metal-Induced Crystallization of Amorphous Ge[END_REF], comparé au processus classique MIC-2 (ALILE) qui procède par l'inversion des couches de Ge et Al[Nast, 2000b].

Table 2 .

 2

1 -Fabrication parameters for the samples imaged in Figure

2

.16. Al and Ge thickness are calibrated using reference layers (not shown).

2.3.2.2 Varying the thickness ratio at constant Ge thickness

		#1 (47B1) #2 (47C) #3 (47H1)
	Layers thickness Al/Ge (nm)	15/20	20/20	40/20
	Ge underlayer (nm)	1	1	1
	Ge interlayer (nm)	2	2	2
	Annealing temperature (°C)	230-260	250	265
	Annealing time (h)	6	12	15
	Interlayer oxidation (min)	9	21	9

Table 2 . 2

 22

-Fabrication parameters for the samples imaged in Figure

2

.17.

Figure 2.17 -Optical images of the samples detailed in

Table 2.2, after annealing. The samples all have the same Ge thickness (20nm) but different Al thicknesses (15 nm, 20 nm, 40 nm).

Table 2 . 3

 23

-Fabrication parameters for the samples imaged in

Figure 2.18.

Table 2 .4 -Fabrication

 2

parameters for the samples shown in

Figure 2.19.

1 (28C1) #2 (28F1) #3 (28I1)

	Layers thickness Al/Ge (nm)	20/24	20/20	20/20
	Ge underlayer (nm)	0	1	2
	AlGeO x interlayer (nm)	2	2	2
	Annealing temperature (°C)	250-300	250-300	240
	Annealing time (h)	20	20	15
	Interlayer oxidation (min)	9	9	9

Table 2 .5 -Fabrication

 2

parameters for the MIC-Ge samples without Ge underlayer (a) and with Ge underlayers of 1 nm (b) and 2 nm (c).

Table 2 .

 2

		(111) GIXRD	(100) L GIXRD
	B (28F2)	0.09	0.12
	C (28E)	0.84	0.74
		(111) GIXRD L	(100)

are indicated by arrows. A 0 is the diffraction of a single al layer. Letter R refers to a reference Ge powder diffractogram. The baseline of each diffractogram is indicated by a solid black line. L 6 -Calculated Lotgering factors L

Table 2

 2

	.7 -Calculated Lotgering factors L (111) GIXRD L	(100)

  -3 ) Thickness and carrier densities of the layers used in the solar cell structure.

	Contact layer (p-GaAs)	300	p-type 1•10 19
	Windows layer (p-AlGaAs)	20	p-type 5•10 18
	Emitter (p-GaAs)	100	p-type 5•10 17
	Base (n-GaAs)	1000	n-type 1•10 17
	BSF (n-AlGaAs)	70	n-type 5•10 18
	Contact layer (n-GaAs)	500	n-type 1•10 19
	Table 4.1 -

Table 4 . 2 -

 42 Table 4.2. IV parameters of the GaAs solar cells characterization and average of the parameters over 7 cells of the sample.

	best cell	0.0056	15.6	0.97	81	12.2
	average	0.0056	15.6	0.79	72	9.2

A (cm 2 ) J sc (mA/cm 2 ) V oc (V) FF (%) Efficiency (%)

Table 4 . 3 -

 43 IV parameters of the solar cell made with 2 µm thick base for the best cell and average parameters over 11 cells (sample 140).

  Table 4.4.A (cm 2 ) J sc (mA/cm 2 ) V oc (V) FF Efficiency (%)

	best cell	0.0056	16.25	0.96 83	12.9
	average	0.0056	16.31	0.95 80	12.5

Table

4

.4 -IV parameter of the solar cells made with 1 µm thick absorber, a higher As 4 :Ga BEP ratio of 27:1 and a higher growth temperature (635-640 °C) (sample 143). Parameter of the best cell and average parameters over 13 cells. No antireflection coating was used in these device.

Table 4 .

 4

							1-R withouth ARC
	(a)	(b)
							EQE without ARC
							1-R with ARC
			1,0				EQE with ARC
	ligth ARC
	light no ARC		0,8
	dark ARC dark no ARC	EQE, 1-Reflectivity	0,4 0,6
			0,2
			0,0
			400	500	600	700	800	900
					Wavelength (nm)
	Figure 4.24 -IV cures with current density in light and dark condition with and without
	70 nm Si 3 N 4 ARC (a) of the best cell (sample #3, 143). EQE and 1-Reflectivity with and
	without the ARC (b).

5 -IV parameter of the final GaAs solar cell on GaAs(111)B substrate, with 70 nm of Si 3 N 4 anti reflection coating (ARC).

  For a Ge powder, reference experimental values are given, normalized to the highest reflection. They are reported in table A.1, as well as the multiplicity determined in section A.1.3.

	2θ(deg) P lane Intensity M ultiplicity
	27.30	111	100	8
	45.34	220	57	12
	53.74	311	39	24
	66.06	400	7	6
	72.88	331	10	24
	83.76	422	17	24
	90.15	511	7	24
	100.86	440	3	12
	107.44	531	11	48
	119.0	620	6	24
	126.65	533	4	24
	141.50	444	2	8
	153.38	711	8	24

Table A.1 -Table of reference diffracted intensities for powder Ge (ICDD card 00-004-0545).

Table A . 3 -

 A3 with an index fix to zero.After correction for the multiplicity, we obtain the value of TableA.3, also normalized to 57 for the 220 reflection. Table of the relative allowed diffracted intensities and corresponding multiplicity for (100)-textured Ge powder in the GIXRD configuration.

	fixed		n°equivalent planes
	0 0 x		2
	0 x x		4
	Permutation depending if the two indexes are equal:
	fixed		n°equivalent planes
	0 k l		2
	0 k k		1
	2θ(deg) P lane Intensity M ultiplicity
	45.34	220	57	4 × 1 = 4
	66.06	400	14	2 × 2 = 4
	100.86	440	3	4 × 1 = 4
	119.0	620	6	4 × 2 = 8
	Inversion (with h and -h index):

http://www.ipvf.fr/

https://en.wikipedia.org/wiki/Solar_energy

Wikipedia:https://en.wikipedia.org/wiki/Differential_interference_contrast_microscopy; Nikon: https://microscopyu.com

https://imagej.nih.gov/ij/ (or https://fiji.sc/)

https://sites.google.com/site/qingzongtseng/template-matching-ij-plugin

http://www.ioffe.ru

that correspond a power of 100 mW. '1.5' is the arctangent of the angle between zenith and 48.2°of latitude and G stand for Global. Global means that it accounts also for the diffuse light.

Pour les croissances de GaAs sur les pseudo-substrats de MIC-Ge, nous avons confirmé que la croissance se fait par épitaxie, mais une forte rugosité augmentant avec l'épaisseur est observée.Nous nous sommes ensuite attelés à la croissance de cellules solaires en GaAs. Une cellule de référence a d'abord été fabriquée et optimisée sur GaAs(111)B, avec une efficacité d'environ ∼16%. Sur les substrats GaAs(111)A et Ge(111), les problèmes de dopage qui apparaissent empêchent d'obtenir des dispositifs viables.Sur les pseudo-substrats MIC-Ge, nous combinons deux problèmes dans la fabrication de la cellule solaire, à savoir la qualité médiocre des cristaux du matériau III-V épitaxié, et l'orientation (111)A des cristaux de GaAs, induisant des difficultés de dopage.Même si l'optimisation future du procédé MIC-Ge pourra améliorer la qualité des cristaux de Ge et celle du matériau III-V associé, la croissance de cellules solaires nécessitera probablement d'autres impuretés dopantes que Si et Be, éléments utilisés dans ce travail.

TEM

Description of the experiment The pole figure experiments may present twin in the GaAs layer. We now want to investigate the possible relationship between the surface pyramids (section 3.4.2.4) and the twinned structure of the layer.

We first study the surface in more detail by SEM and optical microscopy, which allows us to select an area with a particular surface defect for the subsequent crosssectional TEM characterization. We study here a Si-doped GaAs layer, grown using our optimum growth conditions (substrate temperature 600-610°C; As 4 BEP ∼9•10 -6 Torr.) Result Figure 3.51 shows 45°-tilted SEM images of a 450 nm thick GaAs layer. Thanks to a possible channeling contrast, we easily identify the area associated with rotated pyramids (Figure 3.51 (b)).

The sample studied by TEM also comprises an additional 90 nm thick GaAs buffer layer, followed by a thin (Al,Ga) As marker, before the main GaAs layer followed by the standard (Al,Ga)As (3 nm) and GaAs (2 nm) cap layer (Figure 3.52 (b)).

We first select in a Nomarski image twinned with reversed pyramids (white circle in Figure 3.52 (a)). We then prepare a TEM lamella by transversally sectioning the sample with the focused ion beam (FIB) so that it includes both part of the defective and defect-free areas (yellow line in Figure 3.52 (a)).

Using dark field TEM, we select a particular diffraction spot which can discriminate between the two twin variants. In the dark field image, the twinned crystal appears darker than reference crystal (Figure 3.53 (a)) at the position of the rotated pyramids.

In the large area where all the pyramids point in the same direction, we can see that the cross-section TEM shows a perfectly defect-free structure, Figure 3.53 (b).

Chapter 4

Solar cells on GaAs (111)A & B and MIC-Ge substrates 4.2 State of the art of GaAs solar cell on defective Ge layers

We first briefly review the solar cells base on GaAs absorber found in the literature, with a narrow focus on those fabricated on alternative (i.e.defective) Ge substrates.

The idea of replacing the expensive GaAs wafer by a cheaper alternative is not new. GaAs solar cells have already been grown on Ge polycrystalline layers [Venkatasubramanian, 1996;Mauk, 1997] or Ge buffer layers on Si substrates [Wang, 2017], in order to reduce the cost of a simple or tandem III-V solar cells [Hutchby, 1994;Wojtczuk, 1990]. All these approaches take advantage of the close lattice match between Ge and GaAs, which still requires about 1% of In in GaAs to obtain exactly identical cell parameters [Takamoto, 2001]. [Wang, 2017] and (b) Venkatasubramanian et al. on a poly-Ge substrate [Venkatasubramanian, 1996].

Top contact on p-GaAs

Description of the experiment The top and bottom contact are optimized for GaAs p-type and GaAs n-type materials, while on Ge wafers it has to been tested.

In this section we verify that the behavior of the bottom (top) contact on N-type (P-type) GaAs ( 111)B and we compare with a N-type Ge wafer. . To this purpose we made a transmission line measurement (TLM) in order to determine the contact resistance. We use a mask as shown in Figure 4.14 to deposit contact pads at different distance each others.

. For the top contact we deposit a Ti/Au 20/200 nm on a p-type GaAs layer (Be doped layer of ∼0.5 µm on a GaAs (111)B SI substrate). For the back contact, we deposit a Ni/Au/Ge/Au/Ni/Au (4/10/60/110/10/100 nm thick) directly on the n-type GaAs (111)B and n-type Ge (111) wafer back side.

Resistances are measured before and after an annealing at 380-400 °C for 1 min under Ar+H 2 in a RTA oven.

We measure the resistance between each of the metal pads at different distance

Solar cell on GaAs(111)A, Ge(111) and MIC-Ge pseudo-substrates

We now transfer the reference solar cell structure made on GaAs(111)B to other substrate, monocrystalline GaAs(111)A and Ge(111) wafers and finally MIC-Ge pseudo-substrates, to assess the impact on the device performance.

We report in Table 4.6 the previously optimized full solar cell structure (see section 4.4.1.1, p. 162 and section 4.5.2, p. 175). The structure is the same as for the last GaAs(111)B solar cell, except that the GaAs buffer layer is 400 nm instead of 500 nm.

Layer

Thickness (nm) Doping (cm We use a As 4 BEP flux of ∼ 1.1•10 -5 Torr, and growth rate of ∼1.5 Å/s for the Ga, which match the conditions of high V:III ratio (80:1). The substrate temperature is either 620 °C or 640°C. With respect to the GaAs(111)B optimization, the As 4 :Ga ratio is slightly higher and the substrate temperature is lower to help the Si impurities to incorporate as n-type dopant in GaAs(111)A, see section 4.3.1.2. However these small changes may not be sufficient to obtain a complete n-type polarity.

The samples are processed as described section 4.4.2 for contact deposition and etching.

Monocrystalline substrates: GaAs(111)A and Ge(111)

Surface morphology

We first characterize the surface of the cells by SEM and Nomarski microscopy (Figure 4.25). We compare the roughness of the GaAs cells growth on the different substrates.

Compared to the reference solar cell on Ga(111)B, Figure 4.25 (a), the GaAs (111)A structure grown at lower temperature, Figure 4.25 (b), shows a rough surface with triangular pyramids. We observe a SEM contrast in some domain containing

IV characterization

Results

The sample are processed in individual devices and the cells are characterized with the IV curves (Figure 4.26). For each substrate, we compare the obtained cell on GaAs (111)A and Ge(111) substrate to the reference GaAs(111)B device.

reference GaAs( 111

Discussion

The IV characteristics of the devices obtained on GaAs(111)A and Ge(111) substrates are similar and show a clear short-circuit between the top and bottom contacts.

Conclusion

In this section we evaluate the impact of the type of substrate in the solar cell performance optimized on GaAs (111)B substrates. We first used GaAs (111)A and Ge( 111) wafer, all leading to GaAs(111)A layers and we obtain short-circuited devices, despite the similar processing and the possible doping issue.

We initially expected the same problem to occur on the MIC-Ge pseudo substrates, but the extreme surface roughness and the large defect densities cause the fabricated devices to be very resistive.

Appendix A

Texture quantification in grazing incidence X-ray diffraction

In this chapter, we first give general information about X-ray diffraction for a Ge powder and textured Ge crystal. Then, we move to the specifics of grazing incidence X-ray diffraction and how to estimate the crystal texture in this particular configuration.

A.1 Powder X-ray diffraction

In this section we discuss the standard θ-2θ X-ray diffraction configuration and its application to a reference Ge powder sample.

A.1.1 Powder configuration

A standard method to obtain statistical information about the structure of an ensemble of crystal is X-Ray Diffraction (XRD). Independently of the crystal orientation, the Bragg's law determines the angle between the source and the detector at which crystallographic planes can give coherent scattering. Only the planes which satisfy the equation A.1 can diffract.

(A.1) 2d sin(θ) = n λ where:

• n = 1 (simple order diffraction)

• λ is the wavelength of the impinging X-rays → λ CuKα = 1.540 56 Å

• d is the distance between atomic planes, given by equation A.2 for cubic crystals [START_REF] Prince | International Tables for Crystallography Volume C : Mathematical, physical and chemical tables[END_REF].

A.3. Grazing Incident X-ray diffraction For cubic crystals, only the 222 and 444 reflection are considered.

Two extreme cases are possible:

• perfect (111)-texture → p The rescaling is performed as follows :

Grazing Incident X-ray diffraction

When the layer to be analyzed becomes very thin, the amount of diffracting material is too small to obtain satisfying XRD signal values in the Bragg-Bentano geometry.

Alternatively, if we decrease the incident angle to grazing incidence, we illuminate all the sample surface and maintain a good signal for all reflections. The configuration of grazing incidence X-ray diffraction (GIXRD) is discussed in this section.

A.3.1 Configuration

The X-ray grazing incidence configuration uses a very small (fixed) incident angle for the beam of ω=0.3 °with respect to the sample surface (Figure A.4 (a)). In the 2θχ configuration, the X-ray source is fixed, while and the detector rotates in the plane of the sample along the angle 2θχ, as shown from top view in

List of publications and communications

Publications • Pelati, D., Patriarche, G., Mauguin, O., Largeau, L., Brisset, F., Glas, F. and Oehler, F. (2017)

Oral presentations

• E-MRS Spring 2017, "Kinetics and crystal texture improvements on thin germanium layers obtained by aluminium induced crystallization using oxygen plasma"

• E-MRS Spring 2018, "Defect-free epitaxy of GaAs on Ge 111 by MBE"

Posters

• Journées Nationales du Photovoltaïque (JNPV) 2016, "Mechanism of Metal Induced Crystallization of Germanium as pseudo-substrates for solar cells epitaxy"

• 5ème journée scientifique de l'IPVF (April 2016)

• 6ème journée scientifique de l'IPVF (November 2016)

• GDR Pulse 2017, "Ge homo-epitaxy on Ge patches on silica obtained by Metal Induced Crystallization"

• JNPV 2017,"Aluminium induced crystallization of germanium patterns as pseudosubstrates for III-V solar cells"