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Transparent Ohmic Contact for CIGS Solar Cells based on p-type Aluminum Copper Sulfide Material Synthetized by Atomic Layer Deposition. Nathanaelle Schneider, Loraine Duclaux, Muriel Bouttemy, Cathy Bugot, Frederique Donsanti, Arnaud Etcheberry, and Negar Naghavi ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.8b01647 • Publication Date (Web): 13 Nov 2018 Downloaded from http://pubs.acs.org on November 18, 2018

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ACS Applied Energy Materials

Transparent Ohmic Contact for CIGS Solar Cells based on p-type Aluminum Copper Sulfide Material Synthetized by Atomic Layer Deposition. N. Schneider,*,1,2 L. Duclaux,1,2 M. Bouttemy,1,3 C. Bugot,1,2 F. Donsanti,1,4 A. Etcheberry,1,3 N. Naghavi1,2

1 Institut

Photovoltaïque d’Ile-de-France (IPVF), 30 route départementale 128, 91120

Palaiseau, France

2 CNRS,

UMR 9006, Institut Photovoltaïque d’Ile-de-France (IPVF), 30 route

départementale 128, 91120 Palaiseau, France

3 Institut

Lavoisier de Versailles (ILV), Université de Versailles Saint-Quentin en

Yvelines, Université Paris-Saclay, 45 avenue des Etats-Unis, 78035 Versailles Cedex, France

ACS Paragon Plus Environment

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4 EDF

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R&D, IPVF, 30 route départementale 128, 91120 Palaiseau, France

KEYWORDS. atomic layer deposition, copper aluminum sulfide, ternary materials, ptype transparent conducting material, CIGS solar cells


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ABSTRACT. A p-type transparent conducting material was developed to be used as transparent back-contact for the next generation of Cu(In,Ga)Se2 (CIGS) based solar cells. Cu-S:Al-S thin films were synthetized by atomic layer deposition at temperature Tdep = 150°C, by combining CuxS and Al2S3 cycles obtained from copper acetylacetonate (Cu(acac)2), trimethyl aluminum (TMA) and hydrogen sulfide (H2S). Thin films were characterized notably by X-ray diffraction, Scanning Electron Microscopy and UV-vis spectrophotometry. This synthesis strategy allowed a tuning of the film compositions and optoelectronic properties. Films have a high oxygen content and their growth mechanisms as well as their fine chemistries were explored by X-ray Photoelectron Spectroscopy chemical profiling studies. The layers appeared as intermixed copper sulfide and aluminum oxide networks with some Cu-O and Al-S chemical bonds. Finally, best candidates were successfully applied as transparent ohmic contacts with CIGS with contact resistance between 0.1 and 0.7 Ω.cm-2.


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1. INTRODUCTION

Photovoltaic conversion has experienced great changes in the past few years, both in research and at the industry level, thanks to the growing worldwide interest in photovoltaic energy. In this context, breakthrough innovation is extremely important, with more disruptive research to improve existing fields and to prepare for future technological advancements. Among photovoltaic thin film technologies Cu(In,Ga)Se2 (CIGS) is one of the most promising thin-film materials in terms of efficiency with a world record of 22.9%.1

Progress in the fabrication of semitransparent, bifacial, tandem, ultrathin or nanostructured solar cells have enabled new strategies for photon management in several photovoltaic devices. However, because of the lack of appropriate interfacial materials essential to their fabrication, very few CIGS solar cells based on these new technologies have been fabricated up to now.2–7 In all these cells, new structurations are added or integrated to the existing structure to improve the solar cell efficiency. However, to succeed in their integration in thin film solar cells, the most important task is to be able to replace the metallic Mo based back electrodes by a transparent back contact enabling


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the efficient injection of holes at the hetero-structure interface without any absorption of the light. This transparent contact must also be chemically and electrically compatible with the absorber presenting a good band matching and an ohmic contact without detrimental interdiffusion of the species throughout the absorber.8 The most used transparent conducting oxides (TCO) as back-contact for CIGS are n-type ZnO:Al, SnO2:F (FTO) or In2O3:Sn. An ohmic contact between the CIGS and FTO has been previously reported, but a rectifying behavior of the CIGS/ ZnO:Al contact with the formation of a resistive GaOx layer at the interface has been commonly observed.2,5–7 P-type transparent conductors can be interesting candidates as they could act as an electron reflector while helping the formation of an ohmic back contact. Nowadays, the best transparent conductor materials (TCM) are n-type,9–11 and the synthesis of p-type ones with high opto-electrical properties is still a challenge.11–14 NiO was the first p-type TCM reported.15 In 1997, Kawazoe et al. reported CuAlO2 as another p-type TCM with considerable improvement over NiO.16 This opened the development of an entire series of p-type CuMO2 (M = Cr, In, Sc, Y, Ga, B) delafossite TCMs.17 However, neither highconcentration hole doping nor large hole mobility was achieved in these Cu+-based ptype TCMs.18 The use of copper based chalcogen (S, Se and Te) instead of oxygen has


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shown an improvement of the hole mobility because of their higher-lying p orbitals which favors increased mobility over the oxides. Recently, Cu-alloyed ZnS films have shown very interesting TCM properties.19–21 Also, the Cu-III-VI2 with III=Al, Ga and VI=S, Se chalcopyrites with energy band gaps ranging from 1.7 to 3.5 eV can be attractive materials as p-type TCM.22 Among them, CuAlS2 with a wide direct gap of 3.5 eV can be particularly suitable for the use as a transparent p-type back contact in CIGS based solar cells.23,24 However, up to now, only few works have been devoted to the modeling 23,25,26 and the synthesis of this ternary compound as thin films.27–31

To fully succeed in this interfacial design, not only new p-type transparent semiconductors are needed but also an appropriate coating technique is necessary to minimize the recombination centers at interfaces and to limit the diffusion of elements during the deposition process. Atomic layer deposition (ALD) is well-adapted for the preparation of p-type wide band gap semiconductors that would be implemented in various architectures.13 Indeed, by sequentially repeating self-saturated surface reactions, ALD is a suitable deposition technique to control the growth and the properties of uniform and


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conformal materials under relatively soft conditions.32–34 Though less explored than other material class (metals, oxides, …), various ALD-sulfide materials are now available.35 In contrary, ALD-selenide materials, which are also very promising for that target application, are much more rare notably due to the difficulty to access or handle Se sources.36 Ternary materials by ALD, such as CuInS237 or CuCrO238 are usually obtained by combining binary material growth cycles, leading to what is referred as “supercycle”, or by the successive insertion of the different reactants. However, exchange reactions between the adsorbed species on the growing layer and the gas phase species are often observed, which makes the ALD synthesis of more than two element materials very challenging.

The aim of this work is to study the formation of new CIGS/transparent back contact interfaces based on a Cu-Al-S layer to replace the conventional CIGS/Mo contact. As Mo is opaque, it cannot be applied in bifacial solar cells. In our previous work, a first attempt on the growing of CuxAlySz layers at 220°C by ALD has been performed,27 by combining growth cycles of CuxS39 and Al2S340 obtained respectively from Cu(acac)2 (acac =


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acetylacetonate)/H2S and TMA (trimethyl aluminum)/H2S. The amount of {Cu-S} cycles vs {Al-S} cycles has been varied to study the consequences of the insertion of aluminum into the highly conductive Cu-S matrix. Herein, a lower deposition temperature (Tdep = 150°C) has been chosen to operate within the ALD windows of both binary cycles (e. g. {Cu-S}, {Al-S})39,40 and prevent film contamination due to the thermal decomposition of Cu(acac)2.41 Also, this Tdep allows to open the band gap of Cu-Al-S thin films while keeping good electrical properties, and is more adapted for their insertion in CIGS based solar cells. The impact of increasing the number of Al–S cycles, using one or two consecutive Cu–S cycles, on the growth mechanism and the film properties has been investigated. The best compounds in terms of transparency and conductivity have been further characterized by X-ray photoelectron spectroscopy (XPS) fine chemical profiling studies in order to determine their chemical compositions. The growth mechanism of these films has been discussed and finally their ohmicity with CIGS absorbers has been tested.

2. EXPERIMENTAL SECTION

2.1. Material Synthesis.


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The depositions of Cu-S:Al-S layers were conducted in a F-120 ALD reactor (ASM Microchemistry Ltd.), where the two 5 cm x 5 cm substrates are located face to face within a distance of 1 mm, on borosilicate glass and borosilicate glass/Mo/CIGSe substrates. The CIGSe absorber with a thickness of 500 nm was deposited by co-evaporation.2 Prior to deposition, borosilicate glass substrates were washed by sonification in acetone and iso-propanol (2 x 5 min) while borosilicate/Mo/CIGSe substrates were directly used without any pre-treatment. Copper(II) acetylacetonate (Cu(acac)2, 98%, Alfa Aesar), trimethylaluminum (TMA, AlMe3, Optograde, Rohm and Haas), water, and H2S (99.5 %, Air Liquide) were used as copper, aluminum, oxygen and sulfur sources, respectively. All the chemicals were used without further purification. Nitrogen (N2, 99.9999%, Air Liquide) was used as both, carrier and purging gas. TMA, H2S and water were kept at room temperature, Cu(acac)2 was placed in quartz boat inside the furnace and heated at TCu(acac)2 = 130°C. Experiments were performed at Tdep = 150°C and the pressure in the reaction chamber was kept in the range 1 - 5 mbar. A metal sulfide growth cycle, i.e. a Cu-S or a Al-S growth cycle (noted {Cu-S} and {Al-S}) was achieved via the following procedure: metal source pulse / N2 purge / H2S pulse / N2 purge. Typical times were 0.3


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s for pulses, and 0.5 s for purges. For comparison, Cu-S:Al-O thin layers have been prepared by replacing {Al-S} by a {Al-O} cycle, with the program TMA 0.3 s/ N2 0.5 s/ H2O 0.3 s/ N2 0.5 s. The depositions of Cu-S:Al-S and Cu-S:Al-O films were performed with the following programs: [Cu-S:Al-S] = n . (n1 . {Cu-S} + n2 . {Al-S}) with n comprised between 150 and 3000, n1 = 1, 2 and n2 = 1, 2, 3, 5, 10 for typical growth per cycle (GPC) values = 0.18 – 0.75 Å/cycle; and [Cu-S:Al-O] = n . (n3 . {Cu-S} + n4 . {Al-O}) with n comprised between 150 and 3000, n3,n4 = 1, 2, 5, 10; and refered as “Cu-S:Al-S n1:n2“ and “Cu-S:AlO n3:n4”.

2.2. Material Characterization.

All material characterizations were performed on samples deposited on borosilicate glass. Thin-film thicknesses (in 20 - 40 nm range) and morphologies were determined with a Magellan 400L Scanning Electron Microscope (SEM) provided by FEI. For SEM crosssectional images, samples were tilted by 30°. X-ray diffraction (XRD) studies were performed under grazing incidence (GIXRD) conditions with a PanAnalytical Empyrean


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diffractometer using Cu-Kα radiations for crystallinity determination and phase detection. Transmittance and reflectance spectra were obtained using a Perkin Elmer Lambda 900 spectrophotometer with a PELA-1000 integrating sphere. Electrical measurements were performed on 2.25 cm2 samples at room temperature using four-point probe measurement setup and an ECOPIA HMS-3000 Hall effect measurement system with a permanent magnet of 0.5 T. Values of resistivity, carrier concentration, and electron mobility are the average of at least three measurements.

The chemical compositions of films deposited on glass substrates were investigated by XPS with a Thermo Fisher k-Alpha spectrometer using a monochromatic Al-kα source (1486.6 eV) and charge compensation gun is used (except for Cu-S films, evidencing the insulating character of the layer due to Al incorporation). Depth profile measurements were directly performed on ALD deposited layers, and using Ar+ sputtering to remove surface contamination and penetrate into the layers until the substrate (2 keV Ar+, 10 mA beam energy and 30° incident angle, sputtering rate measured for Ta2O5 = 9.5 nm.min1).

XPS composition determination and spectra fitting procedure was done with


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Avantage© software. Chemical environment study was performed on conventional C 1s, O 1s, S 2p and Cu 2p peaks monitored at high energy resolution (CAE 50eV). As Al 2s and Al 2p peaks overlap with the Cu 3s and Cu 3p peaks respectively, a correction of the global peak intensities from their corresponding Cu 3s and Cu 3p contributions was necessary to provide accurate peak area values for the Al 2s and Al 2p quantification. Hence, a specific corrective methodology was developed to overcome this overlapping Al and Cu photopeaks (Section 3.3). The Cu 3s/Al 2s region was considered despite the overlapping, as it presents the advantage to directly bring comparable composition values (equivalent escape depth of photoelectrons). They have similar sensitivity factors (Scofield factors Al 2s: 0.801 and Cu 3s: 0.732) and this allows a direct and easy view of the relative evolution during the profiling of materials. Binding energy (BE) positions and Full-Width-at-Half-Maximum (FWHM) values are given with 0.1 eV and 0.05 eV standard deviations.

2.3. Device Characterization.


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To characterize the ohmicity between the ALD p+interfacial layers and the CIGSe absorber, ALD p+ interfacial layers were directly deposited on borosilicate glass/Mo/CIGSe absorbers.2,3 The devices were then completed by 400 nm sputtered ZnO:Al. Current-Voltage I-V measurements were carried out on 0.1 cm2 cells under dark conditions.


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3. RESULTS 3.1. Film Deposition Strategy.

The Cu-S:Al-S thin films were deposited by combining {Cu-S} and {Al-S} growth cycles, which led to the supercycle [Cu-S:Al-S]. Deposition temperature was set at Tdep = 150°C, as this temperature is within the reported ALD windows of both binary materials, CuxS39 and Al2S3,40 and our previous study has shown Cu-S:Al-S film with good properties could be achieved at that temperature.27 The film thickness was in the range 20 – 35 nm. The amount of {Cu-S} cycles vs {Al-S} cycles has been varied to study the consequences of the insertion of aluminum into the highly conductive Cu-S matrix in terms of growth mechanism, film properties and ohmicity toward CIGSe absorbers.

3.2. Influence of the supercycle on the film properties.

Structural and morphological properties of the films were determined by GIXRD and SEM. Figure 1 presents the XRD patterns of Cu-S:Al-S films for n = 400, with different (n1, n2) values, along with the one of pure CuxS (n2 = 0). As reported previously by our group, ultra-thin CuxS film is a multiphasic compound with a major digenite Cu1.8S (JCPDS 047-


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1748) phase.39 The incorporation of aluminum influences the observed major crystalline phase as well as the size of the crystallite. Indeed, in case of films with n1 = 1, it changes from digenite to chalcocite Cu2S (JCPDS 046-1195) and a gradual amorphization of the film is observed (no visible diffraction peaks from n2 > 3, Figure 1a). Similar observations are made in case of films with n1 = 2. However, the amorphization is postponed and occurs at relatively lower n1/n2 values (Figure 1b).

Figure 1. Influence of the supercycle (n1{Cu-S} + n2{Al-S}) on the Cu-S:Al-S thin film crystal structure with n1 = 1(a), 2 (b).

The morphology of the Cu:S-Al:S thin films as determined by cross-sectional SEM observations confirmed the GIXRD results. To facilitate the characterization, thicker films were deposited by increasing the number of supercycle and some representative examples are presented on Figure


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2. Pure copper sulfide films appear as polycrystalline material (typical crystallite size = 90 – 140 nm). Upon incorporation of {Al-S} cycle, no apparent grains are observable from n2 = 3. This phenomenon is slightly delayed if n1 = 2.

Figure 2. Cross-sectional SEM images of Cu-S (n = 3000), Cu-S:Al-S 1:1 (n = 3000), CuS:Al-S 2:1 (n = 1200) and Cu-S:Al-S 2:5 (n = 1000).

Figure 3 presents the band-gap energy, as determined from Tauc formalism, and the fourpoint probe resistivity of Cu-S:Al-S thin films. Pure copper sulfide film is conductive (ρ =


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10-3 Ω�.cm) but its low band gap value (Eg = 2.2 eV) may limit its transparency. The insertion of Al-S cycle impacts the optoelectronic properties, by increasing both the effective band gap and the resistivity. If n1 = 1, the effective band gap increases almost linearly with the number of {Al-S} cycle, up to Eg = 3.3 eV (1:10). However, the films become very resistive (ρ up to 104 Ω�.cm). If n1 = 2, similar tendencies are observed but are relatively postponed in comparison to the case where n1 = 1. Overall, only two compositions, Cu-S:Al-S 1:1 and 2:5, appear to have the required conductivity property and an effective band gap large enough to be used as transparent ohmic contact with CIGSe. The transmittance and reflection curves and absorption coefficients of thicker CuS:Al-S 1:1 and 2:5 films, along with copper sulfide are presented Figure 4a. Hall measurements have shown that all films are p-doped and the insertion of Al reduces both the

hole

concentration

and

the

mobility

(Figure

4b).

Those

optoelectronic

characterizations confirm that the transparency of pure copper sulfide film is limited by the band gap at low wavelengths and the absorption by the free carriers in the near IR region. These two phenomena are reduced when Al-S cycles are incorporated. For example, Cu-S:Al-S films 1:1 and 2:5 are p-doped, have high charge carrier


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concentrations (> 3x1020 cm-3), and mobility in the range of 0.4 – 1.0 cm²/Vs. In summary, it is possible to tune the optoelectronic properties by varying the {Cu-S}:{Al-S} ratio. Also, only Cu-S:Al-S thin films corresponding to 1:1 and 2:5 ratios seem to have the required transparency (Eg = 2.4 (1:1), 2.3 (2:5) eV, T > 60%) and conductivity to be used as p+ interfacial layer in CIGSe solar cell.

Figure 3. Resistivity (ρ in Ω�.cm) and effective optical band gap (Eg in eV) of Cu-S:Al-S thin films from different ratios with n1 = 1(a), 2 (b).


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Figure 4. Influence of the supercycle ratio on the (a) transmission (full lines, in %), reflection (dotted lines, in %) and absorption coefficient (, in cm-1); (b) hole concentration (in cm-3) and mobility (in cm²/V.s).

3.3. Fine chemical profiling of the films by XPS.

XPS depth profile measurements were performed to determine the surface and in-depth composition and chemical environments of the three films, Cu-S, Cu-S:Al-S 1:1 and 2:5.

For Cu-S film, high energy resolution spectra of the constitutive elements (as determined by the survey spectrum, not shown) Cu 2p, S 2p, C 1s and O 1s, are presented on Figure 5 at the surface and in the layer at different depths (Ar+-sputtering time : 20, 40 and 60 s), before the Cu-S/glass interface crossing. At the surface, C 1s and O 1s peaks are


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mainly representative of carbonaceous contamination, an ultra-thin oxidized film (minor Cuox contribution visible at 933.7 eV BE) on top of the Cu-S layer (Cu-S contribution at 932.4 eV BE) is evidenced by the Cu 2p spectrum, and the S 2p region shows the presence of sulfate compounds (168.0-169.0 eV),42 and S in Cu-S environment (main peak at 161.7 eV). After the first sputtering sequence (20 s), the contamination layer, sulfate compounds and superficial oxide layer disappear (decrease of C and O contents from 17.3 and 2.3 at% to 4.4 and 1.2 at% respectively). Constant Cu 2p (BE = 932.4 eV, FWHM = 1.26 eV) and S 2p (BE = 161.7 eV, FWHM = 1.00 eV) signatures, characteristic of constant Cu(+I) and S(-II) environments in depth are observed,43,44 with higher intensities as the carbonaceous screening has disappeared. Residual carbon and oxygen signals are registered all along the profile at constant levels and positions (284.0 eV, 533.0 eV), which are attributed to the presence of acac ligand fragments inside the film. An average atomic composition of 63.8% Cu, 30.6% S, 1.2% O, and 4.4% C of the Cu-S film can be determined using all the profile levels except the surface one. We note the high homogeneity of the film, and a Cu/S ratio constant and almost equals to 2 (see Supporting Information), which corroborates the major crystalline phase detected using


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GIXRD. The acac fragment incorporation is evidenced by carbon and oxygen presence with a rather constant C/O = 3.7 ratio and constant peak shapes, in agreement with a unique type of fragment trapped during the process.

Figure 5. XPS high energy resolution spectra acquired during depth profiling of the Cu-S sample: (a) Cu 2p, (b) S 2p, (c) C 1s and (d) O 1s.

As previously mentioned, Al 2p and Al 2s peaks overlap with the Cu 3p and Cu 3s preventing a direct determination of the Al content in the layers. To overpass this difficulty, a specific methodology has been developed to study the Cu-S:Al-S layers. The Cu 2p peak is considered in complement to the Cu 3s/Al 2s spectral window to obtain a diagnostic of the chemical environments, easily dissociated, and a peak area reference. Indeed, the determination of the area ratio (Cu 2p/Cu 3s = 0.084) on a standard metallic Cu sample is used as validation criteria for the fitting procedure of the Cu 3s/Al 2s global


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peak. At the surface, the carbonaceous contamination prevents determining a representative Cu 2p/Cu 3s area ratio as attenuation of the peaks intensities due to the carbon contamination screening differs from the low to the high BE regions. Hence, the surface layer will not be considered, as a precise and reliable composition value being only obtained after the removal of the contamination layer. Fixing the Cu 2p/Cu 3s area ratio sets the Al 2s area, and the Cu 3s/Al 2s spectral window is satisfactorily fitted using one contribution for Al (Al-S and/or Al-O) and two for Cu (Cu-S,O and Cu-S). The CuS:Al-S layers survey spectra and fitted high energy resolution spectra of the constitutive elements Cu 3s/Al 2s, S 2p, C 1s on surface and in the bulk can be found in the Supporting Information. The ratios obtained for all the peaks lead to a maximal 4x10-3 standard deviation, which confirms that the separation between Cu 3s and Al 2s contributions is accurately modelled.

The evolution of the Cu 3s-Al 2s and S 2p peaks (normalized in intensity) with depth of Cu-S:Al-S 1:1 and 2:5 films are presented on Figure 6. Longer abrasion times (about 4x) have been necessary for Cu-S:Al-S 2:5 film due to its slightly larger thickness in


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comparison to Cu-S:Al-S 1:1, but above all to its larger hardness. The qualitative examination of the peak evolutions in depth evidences once again the presence of an ultrathin oxidation layer and sulfate –sulfites only at the extreme surface and constant chemical environments Cu(+I) and S(-II) in the whole layers. The overall concentration profiles of Cu-S:Al-S 2:5 and Cu-S:Al-S 1:1 are presented on Figure 7. As expected from the amount of Cu-S and Al-S cycles employed, the Cu-S:Al-S 2:5 is slightly more enriched in aluminum compared to the Cu-S:Al-S1:1. A residual C signal, similar to the one of Cu-S sample in content (about 5 at%) and with similar XPS signature, is observed all along the profile, which is consistent with the incorporation of acac fragments of roughly similar nature. However, a significant concentration of O (20-35 at.%) is observed, which is only partly correlated to the acac ligand fragment (C/O < 0.25 vs 3.7 for Cu-S film), the rest being in the film network. Assessing that similar ligand fragments are incorporated, the corresponding O value (Oligand) can be subtracted and the residual O content (%Ocorr = %O - %Oligand) is mainly attributed to Al oxide, Al presenting a higher sensitivity to O, and Cu(+I) as predominant copper oxide minor phase.45 The global balance of the Cu-S:Al-S layers (Al+Cu)/(Ocorr+S) is determined equal to 1.00 ± 0.05 whatever the depth, and main


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Cu-S and Al-O chemical bonds are identified. The detailed spectra deconvolution (see Supporting Information) also highlights the presence of Cu-O or Cu-(S,O) and Al-S chemical bonds in minor extent ( 104 Ω�.cm) except for the 10:2 composition (ρ = 10-2 Ω�.cm), where very low amount of aluminum was introduced in the chamber. This latter presents effective band gap value and transparency similar to the Cu-S one.

3.5. Ohmic contact with CIGSe absorbers.

The potential of the above ALD p-type layers as alternative transparent back-contacts for CIGSe based solar cells has been explored. For that, the three best p+-layers with similar thickness of 30-40 nm were compared: Cu-S, Cu-S:Al-S 1:1 and Cu-S:Al-S 2:5. The p+ALD layers were deposited on a borosilicate glass/Mo/CIGSe stack, followed by the


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sputtering deposition of a n+-ZnO:Al layer. Figure 8 shows the SEM cross-sectional images of two Mo/CIGSe/ALD p+ interfacial layer/ZnO:Al stacks with Cu-S and Cu-S:AlS 1:1 as interfacial layers. The ALD layers are highly conformal and covering on the CIGSe surface.

Figure 9 compares the I-V characteristics of the Mo/CIGSe/ALD-p+ interfacial layer/ZnO:Al samples with Cu-S, Cu-S:Al-S 1:1 and Cu-S:Al-S 2:5 and of one without any interfacial layer. As the CIGSe/Mo contact has a low resistance, the current voltage (I-V) characteristics of this structure only depends on the resistance of the p+-interfacial layer. Due to the presence of shunt paths some measurement variations were observed, and only the curves exhibiting the highest and lowest resistivity are displayed. While cells without the p+-ALD layer present a non-ohmic characteristic (Figure 9a), the contact has an ohmic character in all three cases with the p+-interfacial layers since the relation between current and voltage is linear. Samples with Cu-S interfacial layer have contact resistances of 0.3 – 1.1 Ω.cm-2, Cu-S:Al-S 2:5 of 12.5 – 630 Ω.cm-2 while the best performing film


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composition seems to be Cu-S:Al-S 1:1 with a contact resistance of 0.1 – 0.7 Ω.cm-2 (see Supporting Information for mean values and standard deviation).

Figure 8. Cross-sectional SEM images of borosilicate glass/Mo/CIGSe/p+ ALD layer/n+ ZnO:Al stack for Cu-S:Al-S 1:1 (left) and pure Cu-S (right).

Figure 9. I-V dark measurements of (a) borosilicate glass/Mo/CIGSe/n+ ZnO:Al stack; and (b) borosilicate glass/Mo/CIGSe/p+ ALD interfacial layer/n+ ZnO:Al stack with p+ ALD interfacial layer = Cu-S (n = 1000, 40 nm), Cu-S:Al-S 1:1 (n = 250, 35 nm) and Cu-


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S:Al-S 2:5 (n = 150, 30 nm). Corresponding contact resistances (in Ω.cm-2) are directly plotted on the graph. Max and min values correspond to the minimum and maximum contact resistance values of a batch of 25 devices.


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4. DISCUSSION

4.1. ALD film growth mechanism.

The ternary materials are obtained using a supercycle strategy, i.e. mixing CuxS and Al2S3 cycles obtained from Cu(acac)2/H2S and TMA/H2S respectively, which have compatible ALD window temperatures.39,40 Hence, supercycles [Cu-S:Al-S] = n . (n1 . {Cu-S} + n2 . {Al-S}) were used and films of various (n, n1, n2) values were deposited at temperature Tdep = 150°C. Film compositions were different from what can be expected by the linear combination of the binary materials, and indicated exchange processes between the Cu-, Al- precursors and the growing layer, as often observed in ALD growth of multinary materials.13,37,46 A wide range of composition and properties could be accessed by tuning both the (n1/n2) ratio and the n1 value, and were compared to pure copper sulfide films. The incorporation of aluminum in CuxS leads to an increase of the film transparency that is due to a larger effective band gap value (improving transparency at low wavelength range) and lower carrier concentration (preventing absorption by free carriers in the NIR region). It is accompanied by an amorphization of the film and a decrease of its


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conductivity. For films with n1 = 2, those phenomena are somehow postponed and occur at relatively lower n1/n2 values. Typically, Cu-rich films (n1/ n2 > 0.2 with n2 = 1) had properties relatively similar to CuxS films, i.e. poly- micro-crystalline, conductive (ρ < 0.5 Ω�.cm), highly p-doped (N = 6.6x1020 cm-3) and with effective band gap energies below 2.7 eV. Only two compositions, Cu-S:Al-S 1:1 and 2:5, were selected as they present a good compromise between transparency and conductivity. Their fine chemistries were also investigated by XPS and compared to pure Cu-S film.

Due to the overlaps between Cu and Al signals, a specific methodology has been developed for the resolution of XPS analyses of Cu-S:Al-S samples. The high O content (> 22 at.%) and the constant Cu/S and Al/Ocorr ratios within the entire layer show that oxygen cannot only originate from the acac fragment and display a reactivity of the layer during the process. Indeed, residual water or oxygen in the ALD reactor is possible as the system operates at moderate vacuum level (at mbar range), which can allow continuous air introduction. The extremely high affinity of aluminum for oxygen implies its almost entire oxidation in the films, as corroborated by the O content in Cu-S:Al-S films compared


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to pure Cu-S layer. Similar oxygen enrichment has been reported in other ALD or PEALD processes of Al2S3, where even the use of O-free precursors and/or the protection with a ZnS capping layer does not prevent the O contamination.40,47,48 A Cu-S/Al-O gradient in depth (Cu-S toward the surface, Al-O toward the substrate) is visible, especially for the 2:5 film. The linearity of its tendency, the gradual shift of the photopeaks toward higher BE and the hardness of the layer are explained by the progressive increase of the insulating character in depth. Even when considering a possible abrasion artifact, a composition fluctuation cannot be ruled out and argues for a transient regime during the first deposition steps, as expected in ALD. The films are thin (around 50 nm), which emphasizes that this gradient is progressively attenuated when reaching the surface.

As the high oxygen content of the Cu-S:Al-S thin films does not prevent them from having the required properties to be used as ohmic contact, and both CuAlO216,49 and CuAlS228–31 materials have been reported as transparent p-type material, we investigated the deposition of Cu-S:Al-O films. However, a strong insulator character was observed due to the aluminum enrichment in these layers. This suggests that a voluntary introduction


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of oxygen (via H2O pulses) impacts the Cu:Al ratio of the films and promotes the formation of Al2O3.

In all films, copper was only in Cu(+I) oxidative state, as no traces of metallic Cu(0) or Cu(+II) were observed. It implies a reduction of the copper precursor during the film growth. From reported experimental43 and computational50 studies of Cu(acac)2 surface chemistry, the reduction of Cu(+II) to Cu(+I) involved in the growth can be assumed to occur during the first half of the ALD cycle.

Finally, the number of consecutive {CuxS} cycles, i.e. the n1 value, strongly influences the properties of the films and postpones the amorphization and the increased resistivity of the film. Atomic compositions of two other Cu-S:Al-S films with the same n1/n2 ratio value (n1/n2 = 0.2), i.e. Cu-S:Al-S 1:5 and 2:10 were determined by XPS and atomic key ratios compared (Table 1). It confirms a higher Cu content (Cu/Al) in Cu-S:Al-S 2:10 films, accompanied by a lower O content (Ocorr/S). Thus, it shows that CuxS grows more easily on itself than on a Al-(S,O)-type surface and it indicates a direct relationship between the film properties and the surface chemistry involved during the ALD growth. Such aspects


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are of prime importance and should be considered in the ALD growth of every multinary materials.

Table 1. Atomic key ratios of Cu-S:Al-S 1:5 and 2:10 thin films determined by XPS. (Cu+Al)/(S+Ocorr)

Cu/Al

Ocorr/S

±0.1

±0.2

±0.2

1:5

0.9

0.7

5.6

2:10

1.1

1.2

2.4

Cu-S:Al-S

4.2. Ohmic contact with CIGSe absorbers.

When reducing the CIGSe absorber thickness down to 500 nm or replacing the Mo by a TCO, the back contact recombination can become an important factor that has to be managed to maintain an efficient carrier collection in the solar cell and the short circuit current.2,4 To overcome this problem, one solution could be to realize a contact that would block the electrons (electron mirror) while having an ohmic behavior with the holes.51,52 Several options are available; one of them being the introduction at the back contact of a very thin interfacial p+ doped semi-conductor layer with an ideal energy band positioning that will act as a perfect electron mirror at the rear side of the solar cell to prevent back contact recombination. While this interfacial layer should act as a perfect electron mirror, it should make an ohmic contact for the holes. Hence, a p+ doping (N  2x1018 cm-3) along with a relatively wide band gap (Eg  2.4 eV) for a good transparency is required. Moreover, this layer should be as thin as possible. While all the interfacial layers tested in this study (Cu:S, Cu-S:Al-S 1:1 and Cu-S:Al-S 2:5) present hole concentrations higher than


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N = 2x1018cm-3, the sample presenting the best compromise between resistivity, transparency and ohmicity remains Cu-S:Al-S 1:1.

The advantage of deposition of this interfacial layer by ALD are multiple: first ALD allows the deposition of a very thin (

Transparent Ohmic Contact for CIGS Solar Cells ... - ACS Publications

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