# Interfacial Chemistry of Cu(In,Ga)S<sub>2</sub> on GaP/Si: Insights for Tandem Solar Cell Design

Eugène Bertin, Éric Gautron, Ludovic Arzel, Léo Choubrac, Antoine Létoublon, Sylvie Harel, Rozenn Bernard, C. Cornet, Maud Jullien, Olivier Durand, et al.

► To cite this version:

Eugène Bertin, Éric Gautron, Ludovic Arzel, Léo Choubrac, Antoine Létoublon, et al. Interfacial Chemistry of Cu(In,Ga)S<sub>2</sub> on GaP/Si: Insights for Tandem Solar Cell Design. physica status solidi (RRL) - Rapid Research Letters (pss RRL), 2025, 19 (8), pp.2500091. 10.1002/pssr.202500091. hal-05095591

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# Interfacial Chemistry of Cu(In,Ga)S₂ on GaP/Si: Insights for Tandem Solar Cell Design

Eugène Bertin,* Éric Gautron, Ludovic Arzel, Léo Choubrac, Antoine Létoublon, Sylvie Harel, Rozenn Bernard, Charles Cornet, Maud Jullien, Olivier Durand, and Nicolas Barreau*

This article exposes an in-depth analysis of the chemical interaction between Cu(In,Ga)S₂ thin-film and GaP/Si(001) pseudo-substrate, in the framework of the development of high-efficiency tandem solar cells on Si. A rich Cu-Ga-P-S chemistry at the Cu(In,Ga)S₂/GaP interface is uncovered, in the case of Cu(In,Ga) S₂ films grown in Cu-excess conditions. The involved chemical reactions are studied and modeled by breaking the system down to its binary constituents, such as Cu₂-xS and GaP and by characterizing the reaction products. A chemical model is proposed to explain the chemical fluctuations observed at Cu(In,Ga)S₂/ GaP interface. It is found that Cu-rich Cu(In,Ga)S₂ decomposes the underlying GaP film. This knowledge of Cu(In,Ga)S₂/GaP interface chemical reactions can benefit the future development of two-terminal monolithic Cu(In,Ga)S₂/Si tandem solar cells.

## 1. Introduction

Wide-bandgap pure-sulfide Cu(In,Ga)S₂ (CIGS) receives significant attention due to its excellent optical properties. The direct bandgap of CIGS can be adjusted from 1.54 eV, as seen in CuInS₂ (CIS), to 2.43 eV in the case of CuGaS₂ (CGS).[1] This variability makes it an attractive option for single-junction

and tandem photovoltaic (PV) applications.[2–5] A practical approach to achieve high efficiency, cost-effective PV could involve the monolithic integration of wide-bandgap CIGS with crystalline Si (c-Si) leading to highly sought-after high-efficiency and low-cost CIGS/Si tandem solar cells.[6,7] Notably, the CuIn₀.₇₅Ga₀.₂₅S₂ chalcopyrite, with a bandgap of 1.7 eV, should make an excellent top cell absorber in a two-terminal configuration. This potential has resulted in efficiencies exceeding 15% for CdS/CIGS[8,9] and Zn(O,S)/CIGS[10] heterojunction solar cells, deposited on Mo-covered soda lime glass (SLG) substrates by coevaporation. However, to apply these advancements to Si-based tandem PV, it is crucial to replace the opaque Mo/SLG substrate with a Si substrate. This change comes with its own set of challenges, including the lack of an optimal selective electrical contact layer between wide-bandgap chalcopyrites and Si[11] and the potential Cu diffusion into the Si bottom cell.[12]

In our prior work,[13,14] we proposed a strategy to circumvent the difficulties outlined earlier. This involves the use of a thin GaP interlayer between the CIGS and the Si for three reasons. Firstly, a thin GaP interlayer allows the fabrication of a high-quality, adherent monocrystalline CIGS film on Si.[14] Secondly, it may inhibit the diffusion of harmful elements like Cu from the CIGS to the Si. Lastly, GaP may serve as an ideal passivating selective contact for CIGS holes. Detailed motivations for fabricating wide-bandgap chalcopyrite on GaP/Si are further explained in ref. [13].

In a previous study,[14,15] we reported the structural differences between Cu-rich and Cu-poor CIGS films coevaporated on GaP/Si(001). The Cu-poor CIGS film consisted of at least three phases; the main crystalline phase was found to be chalcopyrite (CH) Cu(In,Ga)S₂, coexisting with In-rich thiospinel (TS) CuIn₅S₈, and CuAu-ordered (CA) CuInS₂, all sharing epitaxial relationships to each other and the GaP/Si(001) pseudo-substrate. On the contrary, the Cu-rich CIGS film was single-phase chalcopyrite with a lower density of microtwin defects. In both films, the chalcopyrite crystal exclusively grew with its c-axis aligned with the out-of-plane direction Si[001]. In the Cu-excess case, we demonstrated the absence of any secondary phases, including the CuAu-ordered phase, after KCN-etching.

E. Bertin, A. Létoublon, R. Bernard, C. Cornet, M. Jullien, O. Durand
Univ Rennes
INSA Rennes
CNRS
Institut FOTON - UMR 6082
Rennes F-35000, France
E-mail: eugbe@dtu.dk

E. Bertin, É. Gautron, L. Arzel, L. Choubrac, S. Harel, N. Barreau
Nantes Université
IMN
CNRS
Institut des Matériaux de Nantes Jean Rouxel
Nantes F-44000, France
E-mail: nicolas.barreau@cnrs-imn.fr

<figure><img src="image_2.png" /></figure>

The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/pssr.202500091.

© 2025 The Author(s). physica status solidi (RRL) Rapid Research Letters published by Wiley-VCH GmbH. This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

DOI: 10.1002/pssr.202500091
In this article, we study and compare the morphology of the Cu-poor and Cu-rich CIGS films of ref. [14] deposited using a 1-stage coevaporation process on GaP/Si(001) at 570 °C. We also study an additional Cu-rich CIGS film deposited at 450 °C on GaP/Si(001). We comment the morphology of the CIGS/GaP interface. To explain the different morphologies, we investigate the chemical reactions at the CIGS/GaP interface, in the Cu-Ga-In-S-P system. To that end, we deposit thin films of binary Cu<sub>2-x</sub>S, In<sub>2</sub>S<sub>3</sub> and ternary CuIn<sub>5</sub>S<sub>8</sub> and CuInS<sub>2</sub> films on GaP/Si pseudo-substrates and SLG. Then, we perform a structural investigation based on Raman spectroscopy and X-ray diffraction (XRD) to reveal possible reaction products.

**2. Results**

**2.1. Comparison of the Morphology of Cu-Poor and Cu-Rich**
**epi-CIGS**

The SEM images of the cross sectional and surface morphologies of the Cu-poor and Cu-rich samples are presented in **Figure 1**. In both cases, the bulk CIGS exposed by cleavage along Si[110] exhibits features that may be wrongly interpreted as grain boundaries. However, we showed in a previous study that all samples are epitaxial, and do not possess any random angle grain boundaries, which prompts further analysis.

One can see several differences between the Cu-rich and the Cu-poor samples. On the one hand, the Cu-poor sample in Figure 1a exhibits a compact and adherent CIGS film with a

rather smooth surface and a sharp CIGS/GaP interface. Morphological defects are found to roam the cross section, forming in pyramidal shapes in certain areas, which resembles observations made on epi-Cu(In,Ga)Se₂/GaP/Si.[¹³] On the other hand, the Cu-rich sample grown at 570 °C cross-section and surface in Figure 1b,e, reveal a drastically different morphology, with a much rougher surface. Several distinct types of morphological defects can be observed in the cross section and on the surface. (I) Crevices of 200 to 800 nm in width and (II) fractures 20 nm in width are found on the surface. (III) Voids 50–100 nm in size can be detected at the CIGS/GaP interface.

Crevices (I) are typically observed in KCN-etched Cu-rich polycrystalline films[¹⁶,¹⁷] and easily explained by the following model. In Cu-rich conditions, excess Cu is accommodated by the segregation of Cu<sub>2-x</sub>S phase alongside the CH phase. The former is easily etched away by an aqueous KCN treatment. Therefore, regions previously occupied by Cu<sub>2-x</sub>S yield large crevices after etching.

The fractures (II) displayed on Figure 1c only appear in the Cu-rich sample and seem to pave the surface. They seem to run preferentially in the <100> directions of the crystal. The fractures may be related to strain relaxation. Mechanical stress is expected in the case of high mismatch in materials' thermal expansion coefficient (TEC), which are 2.6, 4.65, and 8 × 10<sup>6</sup>/°C for Si, GaP[¹⁸] and CIGS (assuming it is similar to Cu(In,Ga) Se<sub>2</sub>[¹⁹]), respectively. It is plausible that excessive temperature ramping speed during cooling has led to the rupture of the film. However, this raises the question as to why the Cu-poor sample,

<figure><img src="image_3.png" /><figcaption>**Figure 1.** Cross sectional morphology of the CIGS/GaP/Si sample a) Cu-poor and b) KCN-etched Cu-rich epi-CIGS grown on GaP/Si at 570 °C. c) Surface morphology of the KCN-etched Cu-rich epi-CIGS at 570 °C. Close up view around the CIGS/GaP interface for the as-grown Cu-rich samples fabricated at 450 °C and d) 570 °C substrate temperature, prior to KCN etching.</figcaption></figure>
which in principle has the same TEC does not suffer from the
same fractures. The answer may lie in the fact that the Cu-poor
sample is characterized by a phase mixture CH, TS, and CA,[¹⁴]
which may modify how stress due to lattice or TEC mismatch is
accommodated.

In our samples, the voids (III) are only present in Cu-rich sam-
ple grown at 570 °C substrate temperature, as seen in Figure 1c,d.
In the second Cu-rich sample grown at 450 °C, we do not detect
any interface voids. Highly faceted triangular voids on the sub-
strate side have been observed by Lei et al.[20–22] in the case of
Cu-rich epitaxial Cu(In,Ga)Se₂ (CIGSe) grown on GaAs. They
interpret such voids as Kirkendall voids, which result of the
out-diffusion of Ga from the GaAs to the CIGSe layer.
However, in our Cu-rich sample, the voids seem to span across
the interface and do not show clear triangular faceting. This sug-
gests that the voids present in our Cu-rich sample are not
Kirkendall voids.

**2.2. STEM-EDX Analysis of the CIGS/GaP Interface**

To get more insight into the phenomenon taking place at the
CIGS/GaP interface, we perform a STEM-EDX measurement,
as shown in **Figure 2**. By focusing on the HAADF image and
on the P/S EDX map of Figure 2c,d respectively, one can deter-
mine the location of the CIGS/GaP and the GaP/Si interfaces.
Let us, for now, overlook the Ga/In map of Figure 2a and focus
on the Cu EDX map of Figure 2b. A clear Cu-enrichment of the
CIGS/GaP interface, relative to the Cu content in the CIGS layer,
can be seen.

When compiling the results, one can recall that voids at the
CIGS/GaP interface (i) are only present when the growth takes
place in Cu-rich conditions, (ii) at high growth temperatures of
570 °C and (iii) correlate with the presence of a Cu-enrichment of

<figure><img src="image_4.png" /><figcaption>**Figure 2.** STEM-EDX analysis of the CIGS/GaP/Si double interface for the Cu-rich sample grown at 570 °C. EDX-STEM map for a) In and Ga, b) for Cu, and d) for S and P as well as c) corresponding HAADF image.</figcaption></figure>

the interface. Therefore, we hypothesize that interface voids
could be related to a possible chemical reaction involving
Cu₂₋ₓS, at the CIGS/GaP interface. In the following, we test this
assumption by depositing binary and ternary films of the Cu-In-S
chemical system (Cu₂₋ₓS, In₂S₃, CuIn₅S₈, and CuInS₂) onto
GaP/Si pseudo-substrates and on soda lime glass (SLG) for ref-
erence, starting with Cu₂₋ₓS (see Section 2.2 for more details).

**2.3. Interface Reaction of Cu<sub>2</sub>-<sub>x</sub>S and GaP**

2.3.1. Influence of Substrate Temperature on Cu<sub>2</sub>-<sub>x</sub>S Reaction
with GaP

In this section, we clarify the chemistry of the Cu₂₋ₓS/GaP inter-
face by means of Raman spectroscopy. The SEM morphology of
the two Cu₂₋ₓS films grown on GaP/Si at 450 and 570 °C is avail-
able in SI Figure S6 and S7, Supporting Information. In sum-
mary, we find that the deposited film takes the form of
Cu-rich aggregates of about 100 nm in height on the surface.
The deposited films exhibit a very poor wetting of the GaP sur-
face. **Figure 3** displays the Raman spectra of the thin Cu₂₋ₓS
grown on GaP/Si and on SLG. All Raman spectra recorded on
films grown on GaP/Si exhibit a strong Si substrate peak at
520.8 cm⁻¹, of the same intensity. The reference spectrum of the
corresponding GaP/Si(001) platform alone exhibits two peaks at
365 and 402 cm⁻¹ that can be respectively ascribed to the char-
acteristic transverse and longitudinal optical modes of GaP.[²³]
The GaP signal is low due to its small thickness of 80 nm.
The reference spectrum for Cu₂₋ₓS is extracted from the
Cu₂₋ₓS sample grown on SLG at 570 °C and exhibits a strong
A₁ peak at 475 cm⁻¹[²⁴,²⁵] and a minor peak at 265 cm⁻¹.

Let us compare the Cu₂₋ₓS deposited at 570 °C on GaP/Si and
on SLG, as shown on Figure 3. The characteristic GaP and Si
peaks can be detected, demonstrating that the laser probes the
whole Cu₂₋ₓS/GaP/Si stack, and *a fortiori*, the Cu₂₋ₓS/GaP

<figure><img src="image_5.png" /><figcaption>**Figure 3.** Raman spectra of the GaP/Si alone, of the thin Cu₂₋ₓS films deposited at 570 °C on SLG, at 570 and 450 °C on GaP/Si(001). The data are presented with the background stripped, without normalization.</figcaption></figure>
interface. One can notice that the A₁ mode of Cu₂₋ₓS at 475 cm⁻¹ is lower in intensity on GaP/Si, as compared to the reference on SLG. This decrease in the Cu₂₋ₓS signal peak area by a factor of 10 is compatible with a lower volume of Cu₂₋ₓS on GaP/Si than on SLG. More interestingly, additional peaks marked by red arrows at 279, 310, 389, and 425 cm⁻¹ cannot be explained by modes of Cu₂₋ₓS, GaP or Si, which strongly hints that one or several new compounds emerge from the reaction of Cu₂₋ₓS and GaP. The intense peak at 310 cm⁻¹ can be ascribed to CuGaS₂ (CGS, space group I42d).[26] The intensity of this Raman mode may be related to the fact that the bandgap of CGS (E<sub>G</sub> = 2.43 eV[11]) is very close to the laser excitation energy of 2.41 eV; CGS is most likely in resonant Raman scattering conditions. Since no Ga was evaporated during the binary Cu₂₋ₓS film growth, the Ga must originate from the GaP film. Furthermore, the strong mode at 389 cm⁻¹ and the weak mode at 279 cm⁻¹ could be ascribed to the ternary Cu₃PS₄ copper thiophosphate phase (space group Pmn21).[27] Finally, the strong mode at 425 cm⁻¹ may originate from the ternary Cu₇PS₆ copper thiophosphate (space group P2₁3)[28] or the binary CuP₂ copper phosphide (space group P2₁/c).[29] Indeed, both allow intense Raman mode at 425 cm⁻¹. From a thermodynamic standpoint, CuP₂ is only stable in P-rich conditions, that is, an extreme decomposition speed of the GaP.[30] Furthermore, excess S supplied during the process would most likely react with such a compound. Thus, we argue that the peak at 425 cm⁻¹ is most likely due to Cu₇PS₆. Yet, further investigation is necessary to accurately discriminate the two compounds. Since no P has been supplied during Cu₂₋ₓS, the P must originate from the GaP film. This strongly suggests that the Cu₂₋ₓS decomposes the GaP into its elemental components, and reacts with the Ga and P.

The CuGaS₂, Cu₃PS₄, and Cu₇PS₆ Raman peaks are therefore a fingerprint of the reaction of Cu₂₋ₓS and GaP. In Figure 3, we compare the Raman spectra Cu₂₋ₓS deposited at 450 and 570 °C on GaP/Si. One can see that the A₁ mode of Cu₂₋ₓS at 475 cm⁻¹ decreases with the temperature. On the contrary, the fingerprint of the reaction of Cu₂₋ₓS with GaP increases with the temperature. We conclude that a lower volume of Cu₂₋ₓS has reacted with GaP at 450 °C, and lower volume of product has been formed. While the reaction is clearly hindered by the lower temperature of 450 °C, it is not completely prevented.

We conclude from the Raman spectra in Figure 3 of the thin nominally Cu₂₋ₓS/GaP/Si stacks that the Cu₂₋ₓS is consumed and decomposes the GaP. The products of this reaction are most likely CuGaS₂, Cu₇PS₆, Cu₃PS₄, and least likely CuP₂. To confirm that, we employed XRD. However, due to the small thickness and low reaction time, the product compounds did not exhibit a strong enough diffraction signal for phase identification.

### 2.3.2. XRD Study of Cu₂₋ₓS/GaP/Si Stacks

To enhance product volume and crystallite size, a thicker (>100 nm) Cu₂₋ₓS film was deposited at 570 °C on GaP/Si, and a longer high temperature plateau was applied. We first verify by Raman that the same decomposition products have been made in the thicker Cu₂₋ₓS film. In Figure S1, Supporting Information, the Raman spectrum of the thicker Cu₂₋ₓS sample is plotted against the thin Cu₂₋ₓS film of Figure 3, both deposited

on GaP/Si(001). We note that the signal corresponding to the Cu₂₋ₓS phase at 475 cm⁻¹ is completely annihilated for the thicker Cu₂₋ₓS sample. On the contrary, one can see that the signal corresponding to the CuGaS₂ product increases, together with the fingerprint of P-containing products, Cu₃PS₄ and Cu₇PS₆. This clearly shows that a higher volume of GaP has been decomposed by the Cu₂₋ₓS, the latter being the limiting reactant. A X-ray diffractogram in Figure 4a was recorded in the Si[001] longitudinal direction of the substrate, on Cu₂₋ₓS/GaP/Si and on GaP/Si alone. Strong peaks at the 2θ = 32.85 and 69.28° originate from the GaP/Si(001) substrate. An artifact corresponding to the Kβ-Si(004) can be seen at 2θ = 61.70°. In addition to the peaks related to the GaP and Si, additional peaks at 2θ = 29.08, 34.38, and 72.44° are only present in the Cu₂₋ₓS/GaP/Si stack. These peaks are compatible with the theoretical positions of chalcopyrite CuGaS₂ (112), (004), and (008) planes at 2θ = 29.08, 34.21, and 72.06° respectively, based on crystallographic data for CuGaS₂ from ICSD (ICSD code: 258 647).[31] The intensity of the CGS(004) planes relative to the typically intense CGS(112) planes suggests that the CGS crystal has a strong (00 L) texture along the perpendicular of the sample surface.

To elucidate the texture of the CGS, a reciprocal space map is recorded around the GaP(002) peak, and displayed in Figure 4b. One can note the strong GaP(002) reflection. A moderately intense ellipsoid-shaped reflection is visible above the

<figure><img src="image_6.png" /><figcaption>**Figure 4.** a) XRD ω/2θ longitudinal scan in the Si[001] direction of the thick Cu₂₋ₓS/GaP/Si sample and GaP/Si pseudo-substrate alone. b) Reciprocal space map around the GaP(002) reflection. The diffraction signal is represented in reciprocal space in reduced scattering coordinate (Sₓ, Sᵧ) = (Qₓ, Qᵧ)/2π, with Sₓ contained in-plane and Sᵧ orthogonal to the surface. c) Zoom around the GaP(002) peak, after normalization to the Si(004) peak.</figcaption></figure>
GaP(002) and correspond to the CGS(004) reflection. Most of the CGS(004) diffraction signal is condensed in a small region of reciprocal space, demonstrating the strong texture of the corresponding crystallites. Some fraction of the CGS(004) signal is spread in a ring, which may be due to a high degree of mosaic misorientation. Based on a previous work of epitaxial-Cu(In,Ga) S₂, we conclude that an important CGS volume shares an epitaxial relationship with the GaP/Si(001) pseudo-substrate: CGS[100](001)//GaP[100](001)//Si[100](001) (see ref. [14] and SI Section Comments on the epitaxial relationship of CGS on GaP/Si for extensive details). A second ring-like contribution is visible below the GaP(002) reflection, and is attributed to the CGS(112) reflection. The ring-like shape of the CGS(112) reflection demonstrate that some CGS crystallites are randomly orientated. Therefore, the entire CGS volume is mostly composed of epitaxial domains alongside a few randomly orientated domains.

To study the consumption of the GaP, the GaP(002)
diffraction peak has been recorded before and after the deposi-
tion of Cu₂₋ₓS on GaP/Si(002), as displayed on Figure 4c. The
total peak area decreases by a factor 2.6 after Cu₂₋ₓS deposition.
Furthermore, the peak integral breath β increases from 0.22 to
0.30°, which corresponds to crystallite sizes of 38 and 28 nm
respectively, according to the Scherrer equation. We explain this
effect by a reduction in the coherently scattering domain size in
the growth direction of the GaP film, and we interpret it as a
reduction in GaP film thickness. We conclude that a portion
of the GaP is decomposed by the reaction with Cu₂₋ₓS.

To further investigate the products of the reaction of Cu₂₋ₓS with GaP, a 2θ scan is recorded at a constant incidence angle of 8° and displayed in **Figure 5a**. In constant-incidence 2θ scan, the scattering vector changes direction during the measurement, preventing the analysis of texture. However, such scattering geometry allows us to get an insight into randomly orientated crystallites corresponding to the ring-like reflections of Figure 4b. To confirm the presence of CuGaS₂ and Cu₇PS₆, a refinement with the Le Bail method is performed. As a starting point, crystallographic data for Cu₇PS₆ (ICSD code: 22451 – space group P2₁₃) and CuGaS₂ (ICSD code: 258647 – space group I42d) have been extracted from the ICSD database.[³²] The fit result is shown in Figure S5, Supporting Information, and good agreement between the simulated and calculated peaks is found across the whole 2θ range of 15–85°. In Figure 5, we show the refinement result between 2θ = 26 and 36°, where three intense peaks are visible. The contribution of the Cu₇PS₆(300)/(310) reflection to the peak at 2θ = 29.2° is negligible. Therefore, the two peaks at 2θ = 29.2 and 34.36° can be respectively ascribed to the CGS(112) and CGS(004) reflections. The peak at 2θ = 32.2° is ascribed to the (222) reflection of Cu₇PS₆. For CGS, the lattice parameters extracted from the refinement are a = 5.351(3) and c = 10.448(4) Å, close to the crystal data from ICSD of a = 5.351(3)Å and c = 10.477(7) Å. For Cu₇PS₆, the refinement yields a lattice parameter of a = 9.627(4) Å, which is different from the ICSD crystal data of a = 9.668 Å. Differences may be explained by the presence of strain in the film, possibly due the epitaxial nature of CGS.

Apart from the grayed-out region of the diagram excluded
from the fit, the entire diagram can be explained based on

<figure><img src="image_7.png" /><figcaption>**Figure 5.** a) XRD 2θ scan at *ω* = 8° incidence, to match the X-ray beam print to the sample size. Le Bail refinement based on CGS (ICSD code: 258 647 – space group *I*42d) and Cu₇PS₆ (ICSD code: 22 451 – space group P2₁₃). Window in gray excluded from the fit. b) Pole figure at the theoretical Bragg angle Cu₃PS₄(210) plane 2θ = 28.26°.</figcaption></figure>

Cu₇PS₆ and CGS. The Le Bail refinement method clearly dem-
onstrates the presence of polycrystalline CGS and Cu₇PS₆. While
the XRD pattern does not exclude the presence of CuP₂, it rein-
forces the attribution of the Raman peaks at 310 and 425 cm⁻¹ to
CGS and Cu₇PS₆ respectively. In the grayed-out part of the dia-
gram (2θ = 32.6–33.3°), diffraction signal cannot be explained
based on CGS or Cu₇PS₆. We argue that it could be related to
the (006) reflection of hexagonal CuS (ICSD code: 32 105 – space
group P6₃/mmc). However, no clear attribution can be made on
a single diffraction peak.

Interestingly, the Cu₃PS₄ phase, evidenced by Raman spec-
troscopy, does not appear as a polycrystalline phase in the 2θ scan
of Figure 5a, which suggests it is epitaxial. However, it does not
appear in longitudinal of Figure 4a, which would mean that the
epitaxial relationship is complex and does not lead to a diffraction
peak in the Si[001] direction. To detect the Cu₃PS₄ phase, a
pole figure at the theoretical angle 2θ = 28.26° of the most
intense Cu₃PS₄(210) reflections is performed and displayed in
Figure 5b. Four intense reflections are found in the <110> azi-
muth and are attributed to GaP(111) or CGS(112) planes, whose
theoretical 2θ angles are 28.34 and 29.08°, respectively. Around
each of the latter intense reflections, three satellite reflections are
visible, and are attributed to Cu₃PS₄(210) planes. Interestingly, the fact that these reflections appear as narrow points on the pole figure demonstrates that Cu₃PS₄ crystals share an epitaxial relationship with the CGS or GaP. More details on the determination of Cu₃PS₄ texture can be found in Figure S8 and S9, Supporting Information.

## 2.4. Study of Interface Reaction of In₂S₃, CuIn₅S₈ and CuInS₂ with GaP

We have stated earlier that the CIGS/GaP interface of the Cu-poor sample does not exhibit any morphological defect. Earlier, we assumed that the absence of morphological defect may correlate with the lack of reaction pathways of Cu-poor CIGS to chemically react with the GaP. Cu-poor CIGS is composed of chalcopyrite Cu(In,Ga)S₂ alongside In-rich CuIn₅S₈.[24] In₂S₃ binary compound may also be present during growth. To verify that these compounds do not strongly react with the GaP, we study thin films of In₂S₃, CuIn₅S₈, and CuInS₂ deposited on GaP/Si(001) and on SLG for reference. Similarly to the previous study of the Cu₂⁻S/GaP interface, the films are made thin (<100 nm) so that the entire stacks, and a fortiori the film/GaP interface, may be probed by the laser using Raman spectroscopy.

The Raman spectra, the XRD diffractograms and their analysis are presented in detail in Figure S2, S3, and S4, Supporting Information. We do not detect the typical fingerprint of the decomposition of GaP. To explain all the Raman peaks, we do not invoke additional compounds than the ones expected for our different layered systems. We therefore conclude to the absence of a strong reaction of In₂S₃, CuIn₅S₈ and CuInS₂ with GaP. Despite the presence of Cu in the following compounds, we argue that the thiopsinel structure of CuIn₅S₈, as well as chalcopyrite and CuAu structures of CuInS₂ stabilize the Cu in such a way that the reaction of Cu and S with GaP is not favorable.

## 2.5. Thermodynamical Calculations and Decomposition Model of GaP by Cu₂⁻S

In this discussion, we propose a model for the decomposition of GaP by Cu₂⁻S and we thermodynamically evaluate the Gibbs free energy changes ΔJG° of the corresponding reactions. All calculations are performed at 800 K (526.85 °C) if not stated otherwise, and substances are assumed to be in their standard state. For pure substances and binary compounds, standard enthalpies of formation ΔJH°, entropies S° and associated Gibbs free energy of formation ΔJG° are taken from Barin et al.[32] Based on a methodology developed by Scragg et al.[33] we calculated the ΔJG° of CuInS₂ and CuIn₅S₈ with thermodynamical data of refs. [32,34]. For missing data on CuGaS₂, Cu₃PS₄ and Cu₇PS₆, we estimate the ΔJG° with the following method. ΔJH° at 0 K are taken from OQMD,[35] a database that compiles DFT calculations of ΔJH° for a large number of compounds. Based on Ellingham approximation, i.e., assuming no phase change in that temperature range,[35] we argue that ΔJH° at 0 K constitute a good approximation of ΔJH° at 800 K. The entropy S° at 298 K is estimated from the formula unit volume based on the method developed by[36] for ionic salts.[37] Based on Ellingham approximation, i.e.,

assuming constant heat capacity, we argue that the entropy S° at 298 K is a good approximation of the entropy S° at 800 K. The entropy change of the formation reactions ΔJG° can then be calculated from the estimated entropies and the entropies reported in ref. [32] of binaries at 800 K. Thus, the ΔJG° of CuGaS₂, Cu₃PS₄ and Cu₇PS₆ compounds at 800 K is estimated with the expression ΔJG°(800 K) = ΔJH°(0 K) - 800 · ΔJG° (298 K). More details can be found in SI Section *Details on the thermodynamical calculations*.

Ref. [32] reports that the standard Gibbs free energy of formation ΔJG° of GaP and Ga₂S₃ are respectively -64 and -452 kJ.mol⁻¹, at 800 K (526.85 °C). In general terms, Ga is much more stable in S environment than in P environment. We therefore postulate that the decomposition of the GaP by Cu₂⁻S is driven by the affinity of Ga for S, and by the very high stability of the CuGaS₂ chalcopyrite compound. To illustrate that, we evaluate the change in Gibbs free energy ΔJG°R₁ of the hypothetical reaction R1, corresponding to the formation of one mole of CuGaS₂

$$ \frac{1}{2}\text{Cu}_2\text{S} + \text{GaP} + \frac{3}{4}\text{S}_2 \rightarrow \text{CuGaS}_2 + \text{P}_{\text{white}} \quad (1) $$

The phosphorus product is assumed to crystallize in the form of white phosphorus, which by definition has a ΔJG°(800 K) = 0 kJ.mol⁻¹ (standard reference state of P[32]). We calculate a ΔJG°R₁ = -157 kJ.mol⁻¹ for reaction R1, which hints that the strongly negative ΔJG° of CGS alone is sufficient to tip the free energy balance toward the decomposition of GaP.

We propose the following model for the reaction of GaP and Cu₂⁻S. The decomposition reaction is driven by the reaction of the Ga from the GaP with the Cu₂⁻S, to form CuGaS₂. The epi-taxy relationship of CGS with GaP may arise from the fact that CGS probably crystalizes in place of the epi-GaP crystal, as a result of its decomposition. Due to their similar lattice parameters and crystal structures,[14] the reaction occurs without modification of the crystals orientation. The P originally bound to the Ga is therefore free for further reactions. Initially, the P is in the -3 oxidation state in GaP. In Cu₇PS₆, the Cu and S are in their most stable +1 and -2 oxidation states respectively. Within [PS₆]⁻ groups, the P is in its +5 oxidation state. The same holds for Cu₃PS₄ where [PS₄]³⁻ thiophosphate ions impose a +5 oxidation for P.[38] It is clear that the conversion of GaP to Cu₃PS₄ or Cu₇PS₆ requires an oxidation of P by an oxidative agent. The constant evaporation of S during growth may supply sufficient amount of oxidative agent to reduce the P. We therefore argue that the two following competing reactions (R2) and (R3) occur simultaneously

$$ 4\text{Cu}_2\text{S} + \text{GaP} + 2\text{S}_2 \rightarrow \text{CuGaS}_2 + \text{Cu}_7\text{PS}_6 \quad (2) $$

$$ 2\text{Cu}_2\text{S} + \text{GaP} + 2\text{S}_2 \rightarrow \text{CuGaS}_2 + \text{Cu}_3\text{PS}_4 \quad (3) $$

Both reactions R2 and R3 correspond to the decomposition of one mole of GaP to form one mole of CuGaS₂, and we calculate ΔJG°R₂ = -194 and ΔJG°R₃ = -318 kJ.mol⁻¹. Both reactions correspond to a strongly negative ΔJG°. This supports the observation of spontaneous formation of ternaries of the Cu-Ga-P-S system from Cu₂⁻S and GaP binaries. To elucidate the influence of temperature on the spontaneity of the
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