Romaric De Lépinau

Andrea Scaccabarozzi

Gilles Patriarche

Laurent Travers

Stéphane Collin

Andrea Cattoni

Fabrice Oehler

Evidence and control of unintentional As-rich shells in GaAs 1-x P x nanowires

Introduction

Today's photovoltaic market is dominated by Si-wafer based solar cells with 95% share of production [START_REF]ISE @Fraunhofer[END_REF]. The record efficiency for this technology is 26.7% [START_REF] Green | Solar cell efficiency tables (version 52)[END_REF], approaching the maximum theoretical efficiency of 29.4% [START_REF] Richter | Reassessment of the Limiting Efficiency for Crystalline Silicon Solar Cells[END_REF]. Tandem solar cells combining a top III-V semiconductor top cell with a Si bottom solar cell are currently one of the most studied routes to exceed 30% efficiency. Direct epitaxy of III-V on Si using a transition layer to accommodate the different lattice constants remains a challenge [START_REF] Feifel | Direct Growth of III-V/Silicon Triple-Junction Solar Cells With 19[END_REF]. Direct bonding of the III-V top cell on Si cell, either with a 4-terminal or 2-terminal configuration, hold the current state of art efficiencies with 35.9% [START_REF] Essig | Raising the one-sun conversion efficiency of III-V/Si solar cells to 32.8% for two junctions and 35.9% for three junctions[END_REF] and 33.3% [START_REF] Cariou | III-V-on-silicon solar cells reaching 33% photoconversion efficiency in two-terminal configuration[END_REF], respectively. Semiconductor NWs (NWs) have recently emerged as very promising candidates for a new generation of PV devices. The main advantage of NWs over thin films stems from the possibility to grow high structural quality NWs on mismatched substrates thanks to their small footprint that can accommodate the strain without creation of dislocations. NW solar cells have already been reported with efficiencies reaching 17.8% for InP [START_REF] Van Dam | High-Efficiency Nanowire Solar Cells with Omnidirectionally Enhanced Absorption Due to Self-Aligned Indium-Tin-Oxide Mie Scatterers[END_REF], though for this proof-of-concept the NW array was fabricated by top-down lithography and dry etching of a InP substrate. Solar cells based on NW arrays grown by bottom-up techniques such as the vapor-liquid-solid (VLS) growth mechanism have reached efficiencies of 15.3% [START_REF] Aberg | A GaAs Nanowire Array Solar Cell With 15.3% Efficiency at 1[END_REF] for GaAs and 13.8% [START_REF] Wallentin | InP Nanowire Array Solar Cells Achieving 13.8% Efficiency by Exceeding the Ray Optics Limit Science[END_REF] for InP. Compared to the binary NWs studies reported above, ternary NWs offer more potential in term of band gap tuning. In particular, the energy gap of GaAsP NWs can be tuned to 1.7 eV, which provides the best theoretical efficiency for a III-V top cell on silicon tandem [START_REF] Lapierre R R | Theoretical conversion efficiency of a two-junction III-V nanowire on Si solar cell[END_REF]. To date, only a few teams have reported experimental results on solar devices based on GaAsP NWs [START_REF] Wu | Wafer-scale fabrication of self-catalyzed 1.7 eV GaAsP coreshell nanowire photocathode on silicon substrates[END_REF][START_REF] Holm | Surface-passivated GaAsP single-nanowire solar cells exceeding 10% efficiency grown on silicon[END_REF][START_REF] Tchernycheva | Photovoltaic properties of GaAsP core-shell nanowires on Si(001) substrate[END_REF][START_REF] Wood | Fabrication and Characterization of GaAsP Nanowire-on-Silicon Tandem Photovoltaic Cells[END_REF] and a certain number of epitaxial growth difficulties are still to be addressed [START_REF] Zhang | Self-Catalyzed Ternary Core-Shell GaAsP Nanowire Arrays Grown on Patterned Si Substrates by[END_REF][START_REF] Zhang | Growth of Pure Zinc-Blende GaAs(P) Core-Shell Nanowires with Highly Regular Morphology[END_REF][START_REF] Mohseni | GaP/GaAsP/GaP core-multishell nanowire heterostructures on (111) silicon[END_REF][START_REF] Metaferia | GaAsP Nanowires Grown by Aerotaxy[END_REF]. Among others, the alloy homogeneity of ternary NWs is debated. Composition variations between different NWs grown in the same experiment [START_REF] Tchernycheva | Photovoltaic properties of GaAsP core-shell nanowires on Si(001) substrate[END_REF][START_REF] Zhang | Influence of Droplet Size on the Growth of Self-Catalyzed Ternary GaAsP Nanowires[END_REF] have been observed, as well as variations within a single NW along the growth axis [START_REF] Holm | Surface-passivated GaAsP single-nanowire solar cells exceeding 10% efficiency grown on silicon[END_REF][START_REF] Glas | Nucleation Antibunching in Catalyst-Assisted Nanowire Growth[END_REF][START_REF] Mohseni | Structural and optical analysis of GaAsP/GaP core-shell nanowires[END_REF]. Compositional changes have also been observed at the corners of the hexagonal cross-section of AlInP [START_REF] Sköld | Phase Segregation in AlInP Shells on GaAs Nanowires[END_REF][START_REF] Wagner | Growth and segregation of GaAs-AlxIn1-xP core-shell nanowires[END_REF], AlGaAs [START_REF] Rudolph | Spontaneous alloy composition ordering in GaAs-AlGaAs core-shell nanowires[END_REF][START_REF] Mancini | Three-dimensional nanoscale study of Al segregation and quantum dot formation in GaAs/AlGaAs core-shell nanowires[END_REF], and GaAsP [START_REF] Zhang | Polarity-driven quasi-3-fold composition symmetry of self-catalyzed III-V-V ternary core-shell nanowires[END_REF][START_REF] Himwas | In situ passivation of GaAsP nanowires[END_REF] core-shell NWs. These features have been linked to different diffusion lengths for the species during shell growth.

Beside segregation at the edges during radial growth, it is commonly reported that unintentional radial growth occurs concurrently to the VLS axial growth. This has been particularly investigated in binary III-V nanowires, for which the concurrent VLS and VS growth results in particular NW shapes [START_REF] Küpers | Diameter evolution of selective area grown Ga-assisted GaAs nanowires[END_REF][START_REF] Sartel | Effect of arsenic species on the kinetics of GaAs nanowires growth by molecular beam epitaxy[END_REF][START_REF] Colombo | Gaassisted catalyst-free growth mechanism of GaAs nanowires by molecular beam epitaxy[END_REF][START_REF] Gibson | Study of radial growth in patterned self-catalyzed GaAs nanowire arrays by gas source molecular beam epitaxy[END_REF][START_REF] Munshi | Position-Controlled Uniform GaAs Nanowires on Silicon using Nanoimprint Lithography[END_REF][START_REF] Oehler | Measuring and Modeling the Growth Dynamics of Self-Catalyzed GaP Nanowire Arrays[END_REF]. For ternary III-V NWs, the unintentional radial growth creates a core-shell structure, in which the chemical composition of the shell often differs from that of the NW core, as reported for AlGaAs [START_REF] Lim | Controlled growth of ternary alloy nanowires using metalorganic chemical vapor deposition[END_REF][START_REF] Chen | Self-directed growth of AlGaAs core-shell nanowires for visible light applications[END_REF], InGaAs [START_REF] Guo | Phase separation induced by Au catalysts in ternary InGaAs nanowires[END_REF][START_REF] Mohseni | In x Ga 1-x As Nanowire Growth on Graphene: van der Waals Epitaxy Induced Phase Segregation[END_REF], InAsSb [START_REF] Xu | Faceting, composition and crystal phase evolution in III-V antimonide nanowire heterostructures revealed by combining microscopy techniques[END_REF] or GaAsSb [START_REF] Conesa-Boj | Gold-Free Ternary III-V Antimonide Nanowire Arrays on Silicon: Twin-Free down to the First Bilayer[END_REF] NWs. In GaAsP NWs, alloy variations between the unintentional NW shell and the core were observed independently of the growth method, MOCVD [START_REF] Sun | Spontaneous formation of core-shell GaAsP nanowires and their enhanced electrical conductivity[END_REF] or MBE [START_REF] Mohseni | Structural and optical analysis of GaAsP/GaP core-shell nanowires[END_REF][START_REF] Himwas | In situ passivation of GaAsP nanowires[END_REF]. The presence of such unintentional shell with a different chemical composition compared to the NW core poses problems for the design and fabrication of NW-based optoelectronic devices.

For GaAsP nanowires, the presence of the unintentional shell and the factors that influence the composition mismatch between the core and the shell still need to be investigated. In this work, we study the unintentional shell that forms during the axial growth of self-catalyzed GaAsP NWs by using specific growth designs with well-controlled alloy compositions. We accurately measure the core and shell alloy compositions and we estimate the shell thickness using high angle annular dark field (HAADF) imaging and energy dispersive X-ray (EDX) chemical contrast. We also show how the substrate temperature can be used to limit the compositional difference between the NW core and shell.

Method

GaAsP NW ordered arrays were grown on boron-doped Si(111) substrates (resistivity 0.01-0.02 Ohm.cm, from Siltronix) using a patterned silica mask to localize the NW growth. A 25 nm thick silica layer was deposited by PECVD on fully deoxidized Si surfaces (HF 5% for at least 1 min, until the surface is hydrophobic). The silica mask was patterned with a hexagonal array of holes (pitch 500 nm) by electron beam lithography (Vistec EBPG 5000+, 100 KV) using a PMMA resist (495PMMA, 2%). Reactive ion etching based on SF 6 /CHF 3 gasses was used to transfer the pattern into the silica mask. The sample was then cleaned successively in trichloroethylene, acetone and isopropanol. An oxygen plasma was used to remove organic contaminations. Before the introduction of the sample in the MBE load-lock, a HF treatment (HF 1%, 30s) was performed to fully expose the Si(111) surface at the bottom of the holes in the silica mask.

The patterned substrates were then outgazed at 500°C in the MBE system (RIBER 32) until the pressure reading reduces below 10 -9 torr. The GaAsP NW growth was performed with elementary sources, using Ga as the VLS catalyst, in the so-called self-catalyzed growth mode. Gallium, Arsenic and Phosphorous were respectively provided using a standard effusion cell, a valved cracker source producing As 4 tetramers (cracker temperature 600°C) and a valved cracker source producing P 2 dimers (cracker temperature 900°C).

The substrate temperature was monitored by a pyrometer (IRCON 0.9micron) and maintained constant during the NW growth. Depending on the experiment the substrate temperature was fixed between 590°C and 620°C.

All elementary fluxes, including group V, were calibrated atomically using RHEED oscillations on GaAs(100) or GaP(100) surfaces (see references [START_REF] Rudolph | Direct observation of a noncatalytic growth regime for GaAs nanowires[END_REF][START_REF] Ramdani | Arsenic pathways in selfcatalyzed growth of GaAs nanowires[END_REF]). The Ga flux was maintained constant at the equivalent growth rate of 0.16 or 0.18 nm.s -1 .The As and P fluxes were varied so that the atomic flux (As+P):Ga value is between 1.5 and 1.8. The growth was started by opening the Ga shutter only (between 30 sec and 60 sec) at the growth temperature to deposit selectively each silica hole. Then the group V shutters were opened to grow the GaAsP NWs. The Ga cel temperature was the same for the pre by closing all shutters and valves abruptly, and turning off the oven for a rapid cooling.

After the growth, the NWs were dispersed on copper grids and chara electron microscopy (TEM) in a transmission electron microscopy (STEM) observations, bright field (BF) and high angle annular dark field (HAADF) images were collected to chemical analysis acquired using a Brucker Super X detector. Conventional dark field TEM was used to highlight the possible twinning in the NW by selecting a single diffraction spot specific to on the crystalline orientations.

Results

To experimentally evidence the formation of the GaAsP NWs we devise a dedicated growth sequence atomic composition. For this experiment, t shows this specific growth sequence long), continues with a GaAs 1-x P GaAs 1-x P x segment with x≈0.2 (approximately correspond to experimental atomic flux ratio respectively. The actual alloy composition of each N following sections. 3 at the growth temperature to deposit selectively a silica hole. Then the group V shutters were opened to grow the GaAsP NWs. The Ga cel temperature was the same for the pre-deposition step and the NW growth. The growth is terminated by closing all shutters and valves abruptly, and turning off the oven for a rapid cooling.

were dispersed on copper grids and characterized by transmission in a FEI Titan THEMIS tool operating at 200kV. During the scanning transmission electron microscopy (STEM) observations, bright field (BF) and high angle annular dark field (HAADF) images were collected to complement the energy dispersive X-ray spectroscopy (EDXS) chemical analysis acquired using a Brucker Super X detector. Conventional dark field TEM was used to highlight the possible twinning in the NW by selecting a single diffraction spot specific to on To experimentally evidence the formation of the unintentional shell during the growth of ternary GaAsP NWs we devise a dedicated growth sequence, which alternates three segments of different periment, the substrate temperature was fixed to shows this specific growth sequence, which starts with a pure GaP segment ( P x segment with x≈0.5 (approximately 2 µm long) and fi approximately 2 µm long). These approximative alloy compositions correspond to experimental atomic flux ratios ε P = P/(As+P) of ε P =1, ε P = 0. 56 andε . The actual alloy composition of each NW segment will be investigated in detail in the the obtained NW array. The overall yield of vertical nanowires 140nm in diameter to 180 nm, 5 µm long and terminated by a Ga are tilted due to strain from the shell heterostructures. They show no kink or specific defect between the three segments by SEM. Representative 5 µm long a liquid Ga droplet in silica hole. Then the group V shutters were opened to grow the GaAsP NWs. The Ga cell deposition step and the NW growth. The growth is terminated by closing all shutters and valves abruptly, and turning off the oven for a rapid cooling. cterized by transmission operating at 200kV. During the scanning transmission electron microscopy (STEM) observations, bright field (BF) and high angle annular dark ray spectroscopy (EDXS) chemical analysis acquired using a Brucker Super X detector. Conventional dark field TEM was used to highlight the possible twinning in the NW by selecting a single diffraction spot specific to one of shell during the growth of ternary three segments of different he substrate temperature was fixed to 600°C. Figure 1(b) which starts with a pure GaP segment (approximately 1 µm long) and finishes with a long). These approximative alloy compositions = 0.56 and ε P = 0.18, W segment will be investigated in detail in the

: (a) SEM micrograph of the GaAsP NW array. (b) Scheme of the nominal structure of the BF image of the grown images of the lower and upper heterostructures, showing twin defects and contrast changes due to composition h direction obtained by

nanowires is about 50%. The long and terminated by a Ga They show no kink or Representative 5 µm long NWs were selected for TEM studies. Using STEM-HAADF segments. Figures 1(d) and 1(e) detail the lower, ε heterostructures, respectively. number of twins, independently TEM images, Figures 1(d) and 1(e), demonstrate that the are only couple of monolayers thick are relatively well controlled, even without employing al. [START_REF] Priante | Abrupt GaP/GaAs Interfaces in Self-Catalyzed Nanowires[END_REF].

A simple EDX analysis using a line scan along the NW axis, figure 1( three segments of distinct chemical composition However, the reported composition values are rather pure GaP (see Figure 1 . The crystal phase of the NWs is zinc-blende (cubic) with a large of the alloy composition. The sharp contrast variations in each of the , Figures 1(d) and 1(e), demonstrate that the transition regions between two segments are only couple of monolayers thick. This fact indicates that the flux transients in the group V cells even without employing specific flux switching methods using a line scan along the NW axis, figure 1(f), reveals the expected structure: three segments of distinct chemical compositions with increasing As content from bottom to tip. However, the reported composition values are rather unexpected: the bottom ), appears here as GaAs 1-x P x with x ≈ 0.8. This clearly demonstrates that the ) does not capture the actual structure of the NW. improved schematic of the NW structure, in which the axial and radial growth modes have been taken into account to represent the actual alloy compo growth step. The first segment only consists in binary GaP (ε volumes created by the axial and radial growth are of the same composition (light results in the expected GaAs 1-x P x segment by axial growth ( shell is also grown radially (orange). In the final growth step (ε P by axial growth (light red) together with an additional shell (dark the chemical compositions of the NW core and shell grown in same nominal conditions can differ due to the distinct growth mechanisms and flux geometries. resolution chemical map of the lower heterostructure (ε the corresponding HAADF image acquired along a <112> zone axis. This specific orientation provides when investigating the possible presence of a NW shell, as it will be discussed later The Ga elementary map is uniform, as expected, but the As and P composition maps reveal the , figure 1(c), we can clearly distinguish each of the three → 0.56, and the upper, ε P =0.56 → 0.18 blende (cubic) with a large sition. The sharp contrast variations in each of the transition regions between two segments indicates that the flux transients in the group V cells methods like Priante et

), reveals the expected structure: with increasing As content from bottom to tip.

bottom segment, nominally . This clearly demonstrates that the ) does not capture the actual structure of the NW.

(a) Updated scheme the NW structure, including the radial and axial volumes created at multishell structure, depending on the axial position the <112> zone-axis, which reveal a shell structures at the lower and upper heterostructure, poor) compared to the NW core.

schematic of the NW structure, in which the axial and radial alloy composition and material step. The first segment only consists in binary GaP (ε P =1), as the volumes created by the axial and radial growth are of the same composition (light green color). The segment by axial growth (yellow) while = 0.18) the last NW shell (dark red). Note that the chemical compositions of the NW core and shell grown in same nominal conditions can differ due resolution chemical map of the lower heterostructure (ε P = 1 → 0.56) and zone axis. This specific orientation provides , as it will be discussed later. ut the As and P composition maps reveal the presence of a As rich (or P poor) GaA the axial position of the heterostructure. Detailed analysis reveals that this As rich shell actually separates into two GaAs 1-x P x layers, the inner (x in the outer shell, see white lines on the figure 2(b). will be presented later in the discussion section.

Figure 2(c) presents the detailed analysis the lower heterostructure, the Ga map is uniform while the As and P map show opposite variations.

Here we observe an As rich shell the heterostructure in the NW core. In Figure 3, we show a detailed chemical analysis of the lower heterostructure shown We focus here on the chemical composition of across the heterostructure in the NW core. The shells are regularly sampled using rectangular area to obtain a local averaged EDXS chemical composition, see Figure 3(a). Figure 3(b) obtained composition against their position observe a slight decrease of the P content, or an increase of the As, as we move from the lower GaP core (d<0) to the next GaAsP NW core ( presence of a As rich (or P poor) GaAs 1-x P x shell (≈ 20nm thick) around the NW core, irrespectively of the axial position of the heterostructure. Detailed analysis reveals that this As rich shell actually layers, the inner (x≈0.4) and outer shells (x≈0.2) with a larger As conte in the outer shell, see white lines on the figure 2(b). Precise quantifications of the alloy composition in the discussion section. In Figure 3, we show a detailed chemical analysis of the lower heterostructure shown We focus here on the chemical composition of each shell, inner (orange color) and outer (blue color), across the heterostructure in the NW core. The shells are regularly sampled using rectangular area local averaged EDXS chemical composition, see Figure 3(a). Figure 3(b) ed composition against their position d relative to the core heterostructure. For both shell observe a slight decrease of the P content, or an increase of the As, as we move from the lower GaP <0) to the next GaAsP NW core (d>0).

0nm thick) around the NW core, irrespectively of the axial position of the heterostructure. Detailed analysis reveals that this As rich shell actually ) with a larger As content Precise quantifications of the alloy compositions =0.56 → 0.18. Similarly to the lower heterostructure, the Ga map is uniform while the As and P map show opposite variations.

2, wrapped all around

HAADF (<112> zone axis). The are used for the determination by EDX of locally averaged compositions. (b) Composition profile of the outer and inner shells with respect to the distance to the the average compositions to the left and to right of the

In Figure 3, we show a detailed chemical analysis of the lower heterostructure shown in Figure 2(b). each shell, inner (orange color) and outer (blue color), across the heterostructure in the NW core. The shells are regularly sampled using rectangular areas local averaged EDXS chemical composition, see Figure 3(a). Figure 3(b) shows the relative to the core heterostructure. For both shells we observe a slight decrease of the P content, or an increase of the As, as we move from the lower GaP

c,d). The white dotted line marks the interface between the As core. (e-h) Similar analysis for the composition between core and shell

The observed small variations of the chemical composition in the NW shell with that of the NW core suggest that the shell alloy content may be influenced by other factor impinging the NW sidewalls. In Figure 4, we investigate the effect of the substrate temperature on the NW structure and chemical composition. Here, two GaAs step using the same group III (Ga at the equivalent growth rate of 0.18 nm.s (As+P)/Ga = 1.5). The first sample, Figure 4(a 4(e-h), is grown using a higher substrate temperature, 620°C. To account for the axial growth speed reduction with temperature, the growth duration was increased from experiments were performed in succession to minimize any experimental variation. comparable diameter (190 nm and 170 and 4(e), we observe that they still differ in the change in growth duration. They both present a the contrast change in the DF TEM images. In our growth condition temperature (590°C) leads to more defective structure, Figure 4(a), compared the higher growth temperature (620°C), Figure 4(e 4(f-h)) is extracted from an area situated at (respectively 800 nm and 700nm). We clearly observe the presence of The observed small variations of the chemical composition in the NW shell with that of the NW core the shell alloy content may be influenced by other factors than the elementary fluxes impinging the NW sidewalls. In Figure 4, we investigate the effect of the substrate temperature on the NW structure and chemical composition. Here, two GaAs 1-x P x NWs (x≈ 0.2) are grown in a single Ga at the equivalent growth rate of 0.18 nm.s -1 ) The first sample, Figure 4(a-d), is grown at 590°C, while the second sample, Figure a higher substrate temperature, 620°C. To account for the axial growth speed reduction with temperature, the growth duration was increased from 20 min to experiments were performed in succession to minimize any experimental variation.

190 nm and 170 nm below the droplet). Using conventional TEM, Figure 4(a) still differ in length (2800nm at 590°C and 2000 n They both present a cubic structure with several twins, identified by the contrast change in the DF TEM images. In our growth conditions, the lower substrate temperature (590°C) leads to more defective structure, Figure 4(a), compared the higher growth e). Chemical analysis using STEM-EDX and HAADF area situated at approximately the same distance from the NW catalyst 00nm). We clearly observe the presence of an As rich (P po

respectively. (a)(e) TEMin the ZB structure. The of the NW shell grown at and O chemical maps poor) shell and the NW with no visible change in the chemical

The observed small variations of the chemical composition in the NW shell with that of the NW core s than the elementary fluxes impinging the NW sidewalls. In Figure 4, we investigate the effect of the substrate temperature on 0.2) are grown in a single ) and V flux (ε P = 0.24, ), is grown at 590°C, while the second sample, Figure a higher substrate temperature, 620°C. To account for the axial growth speed min to 40 min. The two experiments were performed in succession to minimize any experimental variation. Both NWs have Using conventional TEM, Figure 4(a) nm at 620°C) despite cubic structure with several twins, identified by , the lower substrate temperature (590°C) leads to more defective structure, Figure 4(a), compared the higher growth EDX and HAADF (Figure 4(b-d) and the same distance from the NW catalyst As rich (P poor) shell around the NW core for the sample grown at 590°C, Figure 4( in the EDX map acquired in the same conditions for the sample grown at 620°C, Figure 4( maps of the O content, Figure 4( surface is limited to the first 2 nm and cannot be confused with the Interestingly, the GaAsP shell is separated from the NW core by a samples, independently of the growth temperature 4(f). The chemical map of the sample grown at low temperature, Figure 4(c), shows that this dark HAADF contrast is associated to spatial resolution of this particular EDX map. To better characterize this P-rich interface, we have figure 5(a-b) presents the details (590°C). We clearly distinguish the dark profile (Figure 5 (c)), reveals that the d Figures 5(d-f) show high resolution HAADF images of the NW grown at high temperature (620°C). The dark contrast is also well visible in respectively. The corresponding HAADF profile to the NW outer surface at the top of the NW and that it is buried below an increasingly thicker shell as we move toward the base of the NW.

around the NW core for the sample grown at 590°C, Figure 4(c), and we cannot distinguish any shell in the EDX map acquired in the same conditions for the sample grown at 620°C, Figure 4( Figure 4(d) and Figure 4(h), confirm that the native oxide 2 nm and cannot be confused with the GaAsP shell.

Interestingly, the GaAsP shell is separated from the NW core by a dark line e growth temperature; see dotted white line in Figure 4(b) The chemical map of the sample grown at low temperature, Figure 4(c), shows that this dark a P-rich region. This feature is not visible Figure 4(g) spatial resolution of this particular EDX map. rich interface, we have performed high resolution STEM the details of the core-shell interface of the NW grown at low temperature (590°C). We clearly distinguish the dark line between the NW core and shell. The HAADF intensity profile (Figure 5 (c)), reveals that the dark region is about 7 monolayers thick (1.4 nm). f) show high resolution HAADF images of the NW grown at high temperature (620°C). The in Figures 5(e) and (f), which detail upper and lower part . The corresponding HAADF profiles, Figure 5(g), reveal that the P-rich interface is at the top of the NW and that it is buried below an increasingly thicker shell as we move toward the base of the NW.

), and we cannot distinguish any shell in the EDX map acquired in the same conditions for the sample grown at 620°C, Figure 4 f) show high resolution HAADF images of the NW grown at high temperature (620°C). The upper and lower parts of the NW, rich interface is close at the top of the NW and that it is buried below an increasingly thicker shell

Discussion

The presence of an unintentional shell of different composition than the NW core obfuscates the simple line scan analysis presented in Figure 1(f). The average composition must be carefully deconvolved using the actual hexagonal geometry, core diameter, shell thickness and shell composition to obtain the correct NW core composition (see Supplementary Information and [START_REF] Mohseni | GaP/GaAsP/GaP core-multishell nanowire heterostructures on (111) silicon[END_REF]).

Table 1 summarizes the measured and corrected chemical compositions for the 3-segment GaAsP NW. Note that the corrected alloy composition of the first segment, x=1, now matches the experimental growth conditions, nominally pure GaP. At 600°C (Table 1), the unintentional shell grown is As rich, Δx = -0.03, independently of the P content in the NW core, x core =0.44 or x core =0.18. This difference is even higher for the sample grown at a lower temperature, 590°C (Table 2) with Δx = -0.09 using comparable group V fluxes.

Table 1: Core and shell compositions in the 3 segments of the GaAs 1-X P x NW presented in Figure 1 and 2, measured by STEM-EDX. The measurements were taken at the lower and upper heterostructure, as indicated in Figure 2. The suffix "-APP" refers to the apparent composition, whereas the suffix "-CORR" refers to the corrected composition, which accounts for the surrounding shell. ΔX CORE-SHELL is the composition difference between the core and the shell formed during the same segment growth. Unintentional As-rich shells on MBE-grown GaAsP NWs have been observed on Au-catalyzed NWs [START_REF] Mohseni | Structural and optical analysis of GaAsP/GaP core-shell nanowires[END_REF] and Ga-catalyzed NWs [START_REF] Himwas | In situ passivation of GaAsP nanowires[END_REF]. In addition to the As-rich shell, a P-enrichment at the <112> corners of the shell was also reported by Himwas et al. [START_REF] Himwas | In situ passivation of GaAsP nanowires[END_REF] and Zhang et al. [START_REF] Zhang | Polarity-driven quasi-3-fold composition symmetry of self-catalyzed III-V-V ternary core-shell nanowires[END_REF]. Following Mohseni et al. [START_REF] Mohseni | Structural and optical analysis of GaAsP/GaP core-shell nanowires[END_REF], the core-shell composition mismatch could be attributed to a P-enrichment of the core due to enhanced P adatom migration on the Si(111) surface towards the VLS catalyst. However other experimental studies have shown that group V diffusion is unlikely on the NW sidewall [START_REF] Ramdani | Arsenic pathways in selfcatalyzed growth of GaAs nanowires[END_REF] and that the migration length of atoms on SiOx masks, as used here, is very small [START_REF] Oehler | Measuring and Modeling the Growth Dynamics of Self-Catalyzed GaP Nanowire Arrays[END_REF]. In a first attempt to explain the different alloy contents between the NW core and shell consider the geometry of the MBE in our MBE, the geometrical correction differ slightly between As and P, configuration of the MBE chamber in Figure 6(b) the P content in the the P content in the total group V flux the local arrangement of the group V cell line) than the shell (red dotted line) such a simplistic growth model, o make compositional predictions account [START_REF] Oehler | Measuring and Modeling the Growth Dynamics of Self-Catalyzed GaP Nanowire Arrays[END_REF][START_REF] Gibson | Model of patterned self-assisted nanowire growth[END_REF][START_REF] Kelrich | Shadowing and mask opening effects during selective-area vapor-liquid-solid growth of InP nanowires by metalorganic molecular beam epitaxy[END_REF][START_REF] Dalacu | Selective-area vapour-liquid-solid growth of InP nanowires[END_REF].

The relative P-to-As incorporation at the sidewall surface incorporation relative to P was largely reported for GaAsP layers or gas-source MBE [START_REF] Lapierre R R, Robinson | Group V incorporation in InGaAsP grown on InP by gas source molecular beam epitaxy[END_REF][START_REF] Hou | In situ determination of phosphorus composition in GaAs 1-x P x grown by gas-source molecular beam epitaxy[END_REF]. The P content in the layer phase for group III-limited growth, which conditions at the NW sidewalls are not known, but as a Ga droplet only forms at the top of the wire we can propose a group III limited regime for the shell. In these conditions, the respective kinetics of P and As incorporation on (110) surfaces Keeping geometrical corrections and all impinging fluxes constant, the local unintentional shell composition is found to depend on that of the NW core local content x = 0.46 on the inner shell grown on the GaP core (first segment), which rapidly decreases to x = 0.40 when grown on the GaAsP core (x could be related to a "lattice-pulling" effect, first reported and desc [START_REF] Stringfellow | The importance of lattice mismatch in the growth of GaxIn1-xP epitaxial crystals[END_REF]. Qualitatively, the epitaxy of a mismatched layer adds an elastic strain term to the total free energy. This tends to displace the solid composition to lower this elastic strain, which is balanced by the resulting increase in the chemical free energy. A precise quantificat necessitate to model the complete strain state in the present the scope of this paper. However, we observe a qualitative match, as the local P composition of the shell decreases with that of the NW core, at locations only separated by a few tens of nanometers, Figure 3(a).

In addition to strain-based effect, we measure a strong difference between the unintentional shell grown at low and high temperatures (Figure 4). Compositional variations of GaAsP with temperature have already been observed in planar GaAsP layers with conflicting results. The P content has been either found to decrease [START_REF] Matsushima | Molecular Beam Epitaxy of GaP and GaAs(1-x)P(x)[END_REF] or to increase with the substrate temperature [START_REF] Woodbridge | Structural properties and composition control of GaAsyP1-y grown by MBE on VPE GaAs0.63P0.37 substrates[END_REF]. The possible interpretations are that an increase of the growth temperature favors the evaporation of As from the substrate [START_REF] Woodbridge | Structural properties and composition control of GaAsyP1-y grown by MBE on VPE GaAs0.63P0.37 substrates[END_REF] or the sublimation of As precursor state on the growing surface [START_REF] Nomura | Composition control of GaAsP grown by molecular beam epitaxy[END_REF], leading to larger P contents in the solid. However the actual growth rate, V/III ratio, and molecular form of the precursors can also play a role [START_REF] Nomura | Composition control of GaAsP grown by molecular beam epitaxy[END_REF].

In the NW core, we observe a decrease of the phosphorus content from x = 0.29 at 590°C to x = 0.24 at 620°C. This variation occurs despite the use of identical fluxes and geometry. As the VLS growth is group V limited, it is very sensitive not only to the direct group V fluxes but also to the re-evaporation of group V species from the liquid droplet [START_REF] Glas | Predictive modeling of selfcatalyzed III-V nanowire growth[END_REF]. The final chemical composition of the material is obtained by balancing the impinging and re-evaporation fluxes of group V element in the liquid droplet. The actual balance depends on the thermodynamics parameters of each species and compound, which are different for As, P, GaAs and GaP [START_REF] Ansara | A binary database for III-V compound semiconductor systems[END_REF]. Compared to planar layer growth, which typically operates in the group III limited growth regime, the VLS growth of the NW core operates in a distinct regime for which results extrapolated from planar layer studies may not be valid.

In our growth conditions, the composition of the VLS grown NW core (group V limited) shows a decreasing P content with temperature while the VS-grown shell (possibly group III limited) has an opposite variation. Therefore, there should exist a growth temperature for which the core and shell composition can be made equal. In our experiment, for a group V flux composition ε P = 0.24, a substrate temperature of 620°C leads to a uniform composition x = 0.24 in the core and in the shell. One could perform this optimization for each group V flux composition, ε P , by either lowering or increasing the substrate temperature depending on the initial starting point.

Even in the case of comparable alloy content in the NW core and shell at high temperature growth (620°C), we still observe a dark region by HAADF between the NW core and shell in Figure 5. The dark HAADF contrast suggests a higher P content, which is confirmed in the EDX analysis of Figure 4(c). The position of this local P-rich region correlates with that of the NW core-shell interface, which is buried below an increasingly thicker shell, Figure 5(e-g).

The analysis of the wire grown at high temperature (620°C) shows that the thin interface between the VS-grown shell and the VLS-grown core remains P-rich, even when the growth conditions have been adapted to match the core and shell chemical compositions. This P-rich material can only form just after or during the VLS growth of the top layer of the NW core, before the VS grown shell covers the NW sidewall. Although the exact origin of this P-rich interface is beyond the scope of this article, we can suggest a few alternatives. First, a very fast reaction of the exposed (110) facet with the vapor phase just after it is formed, i.e. P-enrichment after the VLS growth. Second, a local segregation of P at the triple phase line inside the Ga droplet, i.e. P-enrichment during the VLS growth. We note that some sort of segregation at the triple phase line has been reported in the nanowire literature [START_REF] Connell | Identification of an intrinsic source of doping inhomogeneity in vapor-liquid-solid-grown nanowires[END_REF][START_REF] Hilse | Incorporation of the dopants Si and Be into GaAs nanowires[END_REF] but limited to inhomogeneities in dopant concentrations, orders of magnitude below the compositional variations observed here.

Conclusion

We have unambiguously evidenced the formation of a shell during the VLS growth of self-catalyzed GaAsP nanowires by growing successive core alloy compositions. This unintentional shell is always present around the nanowires and results from the radial growth operating on the NW sidewall in the VS mode. The chemical composition of the shell is usually different from that of the NW core due to the distinct growth mechanisms. By adjusting the growth temperature, we find that it is possible to match the chemical composition of this unintentional GaAsP shell with the GaAsP core. Even in these favorable growth conditions, the interface between the NW core and shell still exhibits a local P-enrichment. While the later permits to precisely localize the position of the core shell interface in our GaAsP NWs, it is not clear how this P-rich region is formed. Yet the removal of such inhomogeneities is important for NW-based devices using ternary III-V semi-conductors.

Figure 1 :

 1 Figure 1: (a) SEM micrograph of the GaAsP NW array. (b) Scheme of the nominal structure of the GaAsP NW, showing the three alloy compositions (not to scale). (c) STEM NW viewed along the <110> zone and upper heterostructures, showing twin defects and contrast changes due to composition variations. (f) Chemical composition profile of the NW along the growt using a STEM-EDX linescan.

Figure 1 (

 1 Figure 1(a) shows the obtained NW array. The overall wires are reverse-tapered from droplet. Most NWs are tilted due to specific defect between the three segments by SEM.

:

  (a) SEM micrograph of the GaAsP NW array. (b) Scheme of the nominal structure of the GaAsP NW, showing the three alloy compositions (not to scale). (c) STEM-BF image of the gr NW viewed along the <110> zone-axis. (d), (e) High resolution STEM-HAADF and upper heterostructures, showing twin defects and contrast changes due to composition hemical composition profile of the NW along the growth direction obtained by

  (b)), appears here as GaA simplistic scheme of Figure 1(b) does not capture the actual structure of the NW.

Figure 2 :

 2 Figure 2: (a) Updated scheme the NW structure, including the radial and axial volumes created at each growth step, which results along the wire. (b-c) TEM HAADF and EDX maps acquired double (b) and simple (c) core respectively. The observed shells are always AsIn Figure2(a), we propose an improved growth modes have been taken into account to represent the volume created at each growth volumes created by the axial and radial growth are of the same composition (light next growth step (ε P = 0.56) results i a GaAs 1-x P x shell is also grown segment is formed by axial growth (light the chemical compositions of the NW core and shell grown in same nominal conditions can differ due to the distinct growth mechanisms and flux geometries.

Figure 2 (

 2 Figure 2(b) shows a high-resolution chemical map of the lower heterostructure (ε the corresponding HAADF image acquired along superior details when investigating the possible presence of The Ga elementary map is uniform, as expected, b

( a )

 a Updated scheme the NW structure, including the radial and axial volumes created at each growth step, which results in a core-multishell structure, depending on the axial position c) TEM HAADF and EDX maps acquired along the <112> zone double (b) and simple (c) core-shell structures at the lower and upper heterostructure . The observed shells are always As-rich (P-poor) compared to the NW core.

Figure 3 :

 3 Figure 3: (a) Detailed analysis of the lower heterostructure by STEM marked areas are used for the determination by EDX of locally averaged compositions. (b) Composition profile of the outer and inner shells with respect to the distance to th heterostructure. Dashed lines are heterostructure.

Figure 2 (

 2 Figure2(c) presents the detailed analysis of the upper heterostructure, ε P =0.56 the lower heterostructure, the Ga map is uniform while the As and P map show opposite variations. Here we observe an As rich shell, approximate composition GaAs 1-x P x with x≈0.2 the heterostructure in the NW core.

Figure 4 :

 4 Figure 4: Comparison between NWs grown at 590°C (a DF micrographs along the <110> zone red and blue arrows indicate zones further analyzed by EDX. (b 590°C, observed by STEM-HAADF (b) and (c,d). The white dotted line marks the interface between the As core. (e-h) Similar analysis for the composition between core and shell

  omparison between NWs grown at 590°C (a-d) and 620°C (e-h), respectively DF micrographs along the <110> zone-axis, highlighting one of the twins in the ZB structure. The indicate zones further analyzed by EDX. (b-d) Details of the NW shell HAADF (b) and EDX showing the corresponding P and O . The white dotted line marks the interface between the As-rich (P-poor ) Similar analysis for the NW grown at 620°C, with no visible change in the chemical between core and shell.

Figure 5 :

 5 Figure 5: Details of the core-shell TEM along the <112> zone-axis location indicated by an arrow in (a observations for the NW grown at 620°C toward the NW top (red rectangle of the dark line at the region used to compute the respective the NW top (red line, offset) and

  shell interface. (a) Overview of NW grown at 590°C axis. (b) Atomically resolved STEM-HAADF of the NW shell ion indicated by an arrow in (a). (c) Corresponding HAADF intensity profile. (d the NW grown at 620°C. (e-f) Details of the shell observed by STEM red rectangle) and NW base (green rectangle). Black arrows region used to compute the respective HAADF intensity profile ) and NW base (green line).

  (g). Chemical that the native oxide on the NW outer dark line in HAADF for both Figure 4(b) and Figure The chemical map of the sample grown at low temperature, Figure 4(c), shows that this dark region. This feature is not visible Figure 4(g) due to the low NW grown at 590°C using dark field HAADF of the NW shell at the (c) Corresponding HAADF intensity profile. (d-g) Similar observed by STEM-HAADF ws mark the position profile in (g), toward STEM-HAADF. The the NW grown at low temperature between the NW core and shell. The HAADF intensity about 7 monolayers thick (1.4 nm).

Figure 6 :

 6 Figure 6: (a) Illustration of the different cell inclinations in our MBE chamber, leading to different effective fluxes on the droplet (dashed l lines). (b) Effective proportion of P atoms at the different growth fronts relative to the P proportion in the vapor phase, only considering directly impinging fluxes. Filled circles are our composition measurements by STEM

Table 2 : Core and shell compositions of the GaAs 1-X P x NWs grown at lower and higher temperature, as presented in Figure 4, measured by STEM-EDX. The measurements were taken at the zones indicated in Figure 4(a) and 4(e). The notation is the same as in table 1. GROWTH TEMPERATURE 590°C 620°C X SHELL 0.20 ± 0.01 0.235 ± 0.01

 2

	LOCATION	LOWER HETEROJUNCTION	UPPER HETEROJUNCTION
	SEGMENT	Bottom	Middle	Middle	Top
	X OUTER SHELL	0.20 ± 0.01	0.16 ± 0.01	0.16 ± 0.01	0.14 ± 0.01
	X INNER SHELL-APP	0.41 ± 0.02	0.36 ± 0.02	-	-
	X INNER SHELL-CORR	0.47 ± 0.03	0.41 ± 0.03	-	-
	X CORE-APP	0.75 ± 0.03	0.38 ± 0.02	0.40 ± 0.02	0.17 ± 0.01
	X CORE-CORR	1 ± 0.04	0.44 ± 0.025	0.44 ± 0.03	0.18 ± 0.01
	ΔX CORE-SHELL	-	0.03 ± 0.04	-	0.04 ± 0.02
	X CORE-APP		0.27 ± 0.01	0.24 ± 0.01
	X CORE-CORR		0.29 ± 0.02	0.24 ± 0.02
	ΔX CORE-SHELL		0.09 ± 0.02	0.005 ± 0.02

Acknowledgements

This project has been supported by the French government in the frame of the "Programme d'Investissement d'Avenir" -ANR-IEED-002-01 and by the ANR projects Nanocell (ANR-15-CE05-0026) and Hetonan (ANR-15-CE05-0009). The TEM tool used in this study was funded by the project TEMPOS (ANR-10-EQPX-50).

8 References

In a first attempt to explain the different alloy contents between the NW core and shell he geometry of the MBE chamber. Due to the different incident angles of th corrections for the direct fluxes impinging the droplet and the sidewalls see Figure 6(a). These simple corrections strictly chamber. Assuming the total incorporation of the group V atoms, we P content in the different solids (NW core, sidewall or planar layers) group V flux, taking the planar surface as a reference. I the local arrangement of the group V cells leads to a higher P content in the NW core (blue line), which loosely matches our experimental results (dots). such a simplistic growth model, only based on directly impinging fluxes, is not accurate enough to s if the re-emitted fluxes and shadowing effect

As incorporation at the sidewall surface must also be considered. relative to P was largely reported for GaAsP layers grown by solid-

The P content in the layer is found lower than the P content in the vapor limited growth, which is justified by kinetics arguments [START_REF] Woodbridge | Structural properties and composition control of GaAsyP1-y grown by MBE on VPE GaAs0.63P0.37 substrates[END_REF][START_REF] Nomura | Composition control of GaAsP grown by molecular beam epitaxy[END_REF] conditions at the NW sidewalls are not known, but as a Ga droplet only forms at the top of the wire we can propose a group III limited regime for the shell. In these conditions, the respective kinetics of on (110) surfaces could also play an important role in the final composition Keeping geometrical corrections and all impinging fluxes constant, the local unintentional shell composition is found to depend on that of the NW core. Using Figure 3(b), we measure a on the inner shell grown on the GaP core (first segment), which rapidly when grown on the GaAsP core (x core =0.44, second segment) pulling" effect, first reported and described by Stringfellow on GaInP he epitaxy of a mismatched layer adds an elastic strain term to the total free energy. This tends to displace the solid composition to lower this elastic strain, which is balanced by the resulting increase in the chemical free energy. A precise quantification of this effect would necessitate to model the complete strain state in the present core-shell NWs [56] the scope of this paper. However, we observe a qualitative match, as the local P composition of the In a first attempt to explain the different alloy contents between the NW core and shell, we first Due to the different incident angles of the P and As cells for the direct fluxes impinging the droplet and the sidewalls strictly depend on the local total incorporation of the group V atoms, we plot s (NW core, sidewall or planar layers) as a function of reference. In our configuration, to a higher P content in the NW core (blue dashed , which loosely matches our experimental results (dots). However nly based on directly impinging fluxes, is not accurate enough to and shadowing effects are not taken into considered. A preferred As -source MBE [47-52] P content in the vapor [START_REF] Woodbridge | Structural properties and composition control of GaAsyP1-y grown by MBE on VPE GaAs0.63P0.37 substrates[END_REF][START_REF] Nomura | Composition control of GaAsP grown by molecular beam epitaxy[END_REF]. The local growth conditions at the NW sidewalls are not known, but as a Ga droplet only forms at the top of the wire we can propose a group III limited regime for the shell. In these conditions, the respective kinetics of the final composition.

Keeping geometrical corrections and all impinging fluxes constant, the local unintentional shell Using Figure 3(b), we measure an averaged on the inner shell grown on the GaP core (first segment), which rapidly =0.44, second segment). This observation ribed by Stringfellow on GaInP he epitaxy of a mismatched layer adds an elastic strain term to the total free energy. This tends to displace the solid composition to lower this elastic strain, which is balanced by ion of this effect would [56], which is beyond the scope of this paper. However, we observe a qualitative match, as the local P composition of the