Effect of sodium diffusion on the properties of CIGS solar absorbers prepared…

Physics-based electrical modelling of CIGS thin-film photovoltaic modules for system-level energy yield simulations

Copper indium gallium selenide (CIGS) is a commercialized, high-efficiency thin-film photovoltaic (PV) technology. The state-of-the-art energy yield models for this technology have a significant normalized root mean square error (nRMSE) on power estimation: De Soto model—26.7%; PVsyst model—12%. In this work, we propose a physics-based electrical model for CIGS technology which can be used for system-level energy yield simulations by people across the PV value chain. The model was developed by considering models of significant electrical current pathways from literature and adapting it for the system-level simulation. We improved it further by incorporating temperature and irradiance dependence of parameters through characterisation at various operating conditions. We also devised a module level, non-destructive characterization strategy based on readily available measurement equipment to obtain the model parameters. The model was validated using the measurements from multiple commercial modules and has a significantly lower power estimation nRMSE of 1.2%.

Introduction

Photovoltaics has evolved from an additional power source in small calculators to a clean energy source in mainstream energy production. The growing concerns over climate change have fuelled further evolution towards innovative concepts like vehicle-integrated PV, net-zero energy buildings and net-zero energy districts. In many of these applications, thin-film PV technologies are more suitable than crystalline silicon (c-Si) solar cells. The thin-film technologies have existed for decades, but c-Si PV has dominated the PV industry with its high efficiency, stability, and mature manufacturing processes 1. The thin-film PV technologies are known for their aesthetic appeal and the possibility to fabricate them on flexible substrates. With such characteristics, thin-film PV technologies may gain market share in the domain of integrated photovoltaics (IPV). A report by Becquerel Institute 2 states that, in 2020, the European IPV market was valued at almost 600 million euros and projected to triple reaching 1800 million euros by 2023. The case study in the report highlights the cost-effectiveness of thin-film technologies. The report also identifies digitization as one of the major factors needed to boost the economy around IPV. A modelling infrastructure for reliable energy yield prediction is a step in that direction as it assists in developing and optimizing thin-film PV systems for maximum energy yield. In this work, we FOCUS on creating an energy yield model for system-level simulation for the CIGS technology.

CIGS solar cells have additional current mechanisms that make the current–voltage (IV) characteristics unique compared to conventional solar cells. In c-Si solar cells, the current–voltage (IV) characteristics can be constructed by super positioning the dark behaviour and the photogenerated current at short circuit. This superposition fails in CIGS solar cells because of the voltage-dependent photogeneration 3,4,5. This can be seen in the measured dark and illuminated IV curves shown in Fig. 1. The current mechanisms that contribute to the photocurrent in CIGS solar cells are drift current across the depletion region to the heterojunction interface, thermionic emission across the junction, bulk recombination, and diffusion to the back contact. The voltage dependence is caused by the drift and the thermionic emission losses that are linked to the change in the electric field and increased back contact recombination 6,7,8. The electric field and the depletion layer width, in turn, are voltage dependent. The parasitic currents in CIGS solar cells are junction recombination current, ohmic shunt current, space charge limited current (SCLC), and tunnelling currents 9. The recombination current in CIGS solar cells varies with illumination as some defects are activated upon illumination 10,11. Shunt current is related to the presence of pinholes in the absorber. The absence of buffer and window layers results in a local metal-semiconductor-metal contact which leads to SCLC 12,13,14,15. The tunnelling currents are caused by mid-gap defects that are significant at temperatures below 250 K 11 and can be ignored for the energy yield estimation.

One of the commonly used models for PV module performance is the Sandia PV Array Performance Model (SAPM) 16. This model uses IV characteristics to estimate fitting coefficients to represent the relation of the IV parameters at different operating conditions. This model was first developed for c-Si modules but has also been used for thin-film modules based on CIGS and cadmium telluride (CdTe) technologies 17 ,18. A database of the fitting coefficients for different modules is publicly available and integrated into open-source software packages such as System Advisory Model (SAM) 19 and pvlib 20. and commercial software like PVsyst 21. The Loss Factor Model (LFM) 22 like SAPM, uses the outdoor measurement data to obtain a set of coefficients to correct temperature and spectral mismatch. LFM and SAPM are both empirical models. These models require a significant amount of onsite measurement data to calibrate the parameters for realistic energy yield estimation.

Physics-based models consider the variation of the physical parameters with the external operating conditions to simulate the performance of solar cells. The availability of such models is minimal. The De Soto model 23 (aka the five-parameter model) is a physics-based model based on the superposition of the dark and the illuminated IV curve of a solar cell. It has been widely used by the PV industry to represent the performance of the c-Si solar cells. PVsyst 21 uses a modified version of the De Soto model where it considers an exponential relation between shunt and illumination intensity instead of a linear relation. The efficacy of De Soto and PVsyst models are tested and reported in the section “Statistical significance”. The analytical model described by Sun et al. 6 is a physics-based model developed specifically for CIGS. It addresses the superposition failure by modelling the voltage dependence of photogeneration. However, this model does not consider the variation of saturation current and shunt resistance with illumination intensity, which are significant at low irradiance. The model was only validated at intensities above 400 W m −2 with laboratory-scale cells. The temperature and illumination dependence validated at the cell level may not be valid at the module level.

The reference parameters are parameters that represent the module characteristics at the Standard Test Conditions (STC). They are inputs to a model to estimate module performance at other operating conditions. Multiple techniques are available for estimating the value of such parameters. The nonlinear least-squares method is the most common technique to fit a model to data and obtain the parameters. It iteratively minimises the least square error. The requirement of initial values and bounded solution space for the parameters become a constraint for this method. Wrong initial estimates may lead to non-convergence or convergence to local minima. As the number of parameters increases, uncertainty in parameter estimation increases. This limits its use in large models. Another method is to reduce the number of model parameters by representing the possible dependent variables as a function of independent variables and iteratively solving the model to correctly estimate the power ( \(P_\mathrm\) ) at the maximum power point (MPP), the open-circuit voltage ( \(V_\mathrm\) ) and the short-circuit current ( \(I_\mathrm\) ) 24. Both the above techniques are suitable for a simple model like the De Soto model.

The model described by Sun et al. 6 has many parameters and they simplify the problem by individually fitting the current curves with their respective model i.e., fitting the reverse-biased dark current to their shunt current model, forward bias dark current to their diode current model. The photocurrent curve is obtained by separating the diode current and shunt current from the illuminated IV curve. Then, the photocurrent model is fitted to the photocurrent curve. For certain parameters, estimates were obtained through the numerical simulation software ADEPT, which requires material-specific parameters (e.g., layer thickness, mobility, doping densities, and defects). This methodology by Sun et al. may be an interesting approach as it reduces uncertainty by splitting the model into three smaller parts. However, this strategy cannot be applied to parameterise the model for a commercial module. It is neither possible to measure the reverse-bias IV curve of commercial PV modules because of the presence of bypass diodes nor to have material-specific data like doping concentration, mobility, etc.

To set a benchmark, the accuracy of the existing state-of-the-art models to estimate the IV characteristics of CIGS modules were analysed. Although the CIGS-specific model with a parameter extraction methodology described by Sun et al. 6 should be considered as the state of the art, it cannot be used at the module level. In this work, the variations of both De Soto and PVsyst models are considered as the state of the art for module-level simulations. The De Soto model is used along with the parameter extraction technique described by Laudani et al. 24. The performance estimation data of the PVsyst model was obtained from the module database within the PVsyst software. The IV curves for different modules were estimated using both models and compared with the measured curves to calculate their accuracy.

Figure 2a compares the Power–Voltage (P–V) curves estimated by the De Soto and PVsyst models with the measured IV curves for different illumination intensities at 25 °C. The De Soto model overestimates \(V_\mathrm\) and \(P_\mathrm\) at low intensities and PVsyst overestimates them at high intensities. Figure 2b compares the model estimates with measured curves for different cell temperatures at 1000 W m −2. At high temperatures, the De Soto model underestimates \(V_\mathrm\) and PVsyst overestimates \(P_\mathrm\). The nRMSE in estimating the maximum power point using the De Soto and PVsyst model are 26.7% and 12%, respectively (The graph representing this data is discussed later in the text).

In this work, we develop an electrical model for energy yield estimation adapting the models of significant current pathways from literature and incorporating the temperature and irradiance dependence of various parameters. These relations were determined by characterization. We also develop a step-by-step characterization strategy to obtain the parameters for using the model. The proposed characterization methods are non-destructive and can be easily done with basic PV lab instrumentation. Apart from model parameters, module temperature and incident irradiance are the only external parameters required for energy yield prediction. Thus, the proposed model can be easily combined with a thermal model and irradiance model in any of the existing energy yield prediction infrastructure. With modular and simplistic approach proposed, we ensure that the model can be used across the PV value chain. Table 1 gives the summary of the state-of-the-art models available for simulating the performance of CIGS devices and compares it with the developments made in this work.

Shunt resistance

Shunt resistance ( \(R_\mathrm\) ) can be determined as the inverse slope of the IV curve at \(I_\mathrm\) ( \(V\) = 0) 27. Assuming same shunt resistance for all cells, single cell shunt resistance can be obtained by multiplying the \(\frac

\) to the module shunt resistance. The measurement of shunt resistance at different incidence irradiance shows that it varies non-linearly with irradiance. We propose to use Eq. (7) to represent the shunt resistance variation for CIGS technology. \(R_\mathrm\_\mathrm\) is the shunt resistance at reference irradiance \((G_\mathrm)\) of 1000 W m −2 and \(G\) is the incident irradiance:

Figure 4a shows the irradiance dependence of shunt resistance for different CIGS modules. The power relation (7) was used to fit the data for each module to obtain the empirical parameter \(u\) that defines the relation. The values extracted from the data are given in inset table of Fig. 4a.

Ideality factor and saturation current pre-factor

The ideality factor (n) and dark saturation current pre-factor \((I_)\) can be estimated from the \(I_\mathrmV_\mathrm\) curve obtained by plotting \(I_\mathrmV_\mathrm\) (cell parameters) at different irradiance intensities. This method is used with an assumption that at \(V_\mathrm\). diode current is equal to photocurrent and shunt current is negligible. However, our model considers voltage-dependent photogeneration, illumination-dependent dark saturation current, and non-negligible shunt current. The value of n and \(I_\) cannot be directly used as the reference parameters. Equation (8) was used to fit the \(I_\mathrmV_\mathrm\) to get an initial estimate for \(n\) and \(I_\).

Curve fitting

At this point, the reference values \(R_\mathrm,\mathrm\). \(E_\mathrm\). and \(R_\mathrm,\mathrm\) (cell parameters) have been estimated. Now, the IV curve at STC can be fitted to the Eq. (5) using python SciPy curve fit function 28 to obtain the remaining parameters \(\alpha _\mathrm\). \(V_\mathrm,\mathrm\). \(n\). \(I_00,\mathrm\). and \(\gamma _\mathrm\). Equations (9) and (10), adapted from 6 give the temperature relation of the parameters \(\alpha\) and \(V_\mathrm\). respectively. Here \(\Delta E_\mathrm\) refers to the conduction Band offset between the buffer and absorber. Value of \(\Delta E_\mathrm\) = 0.1 eV fits well for the CIGS technology. For some modules, γ exhibited slight temperature dependence given by Eq. (11), where m is either 0 or 1:

Effect of sodium diffusion on the properties of CIGS solar absorbers prepared using elemental Se in a two-step process

The influence of Na diffusion from various glass substrates during a high-temperature slenization process on the microstructure and morphology of two-step formed CIGS absorber layers is investigated. In order to minimise the CIGS absorber formation time, elemental Se vapour is used to prepare the CIGS absorber. The grain sizes of the CIGS films are found to increase with increasing sodium in the glass substrates (extra clear glass, soda-lime glass, borosilicate glass). TiN and SiN thin films are used as diffusion barrier layers inserted between the glass substrate and the Mo rear conatct to tune the Na diffusion from the soda-lime glass. The interdiffusion between the In-rich CuInSe2 surface layer and the Ga-rich CuGaSe2 layer is promoted by the barrier layer, leading to larger CIGS grains. Efforts are also taken to understand the differences in Na diffusion (from the glass substrates) and their effects on the MoSe2 intermediate layer formation during the high-temperature CIGS absorber formation processes. We find that excess amounts of Na and Se are essential for the MoSe2 growth. The excessive Na in the form of Na2Sex at the CIGS/Mo interface works as a Se source and catalyses the MoSe2 formation. The Se flow in the two-step CIGS formation process must be sufficiently high to obtain high-efficiency CIGS solar cells.

Introduction

During the two-step CIGS formation process, Na diffusion from the substrate plays an important role in forming a uniform CIGS film: the Na content can postpone the inter-diffusion of In and Ga atoms to form a uniform CIGS layer 15,16,17,18,19. However, the Na diffusion is related to the properties of the rear contact and the selenization process. The rear contact properties have a strong influence on the formation of the intermediate MoSe2 layer and the properties of the formed CIGS absorber layer 12,20,21. The MoSe2 layer plays a significant role in ensuring a low contact resistance and good adhesion of the CIGS absorber to the rear contact. The CIGS layer tends to delaminate from the Mo coated substrate after CdS buffer layer deposition if its adhesion to the substrate is not promoted by the MoSe2 layer, which is normally formed at the Mo and CIGS interface during high temperature CIGS formation process. The existence of a thin MoSe2 layer can decrease the apparent Schottky barrier height and thereby provide a better ohmic contact to the CIGS absorber. Simultaneously, MoSe2 is able to enhance the mechanical peel strength of CIGS to Mo back contact. However, an excessive thickness of MoSe2 layer formed at the Mo and CIGS interface is reported to deteriorate the performance of the CIGS solar cells due to the high resistivity of MoSe2 (10 1 –10 4 Ohm-cm) 14. Formation of MoSe2 depends mainly on the selenisation process and alkali Na diffusion from the soda-lime glass (SLG) substrate. over, the preferred orientation of the MoSe2 crystal grains is important for the adhesion of the CIGS layer to the rear contact. Hence, controlling the MoSe2 layer formation and thickness are important for preparing high-efficiency CIGS solar cells.

In this paper, the selenisation of the metallic Cu-In-Ga precursor was carried out in nitrogen atmosphere at ambient pressure using elemental Se vapour by an inline Rapid thermal processing furnace (Smit Thermal Solutions, Netherlands). There are many advantages in using elemental Se vapour instead of toxic H2Se gas as the Se source, such as lower cost, easier handling and shorter absorber process time of 2Se gas). In this work, efforts are taken to understand the difference in Na diffusion (from the SLG substrate) and their effects on MoSe2 intermediate layer formation during the high-temperature CIGS absorber formation process. Bilayer Mo rear contacts are deposited on glass substrates containing different sodium content.

Experimental Details

The different glass substrates and rear contact designs utilized in this study for CIGS solar cell fabrication are shown in Fig. 1. The optimized deposition conditions of the individual layers used to form various back contact stack designs are discribed in our previous publication 18. Bilayer Mo films comprising of a high pressure (HP, 6.0 × 10 −3 mbar) deposited bottom Mo layer and a low pressure (LP, 1.5 × 10 −3 mbar) deposited top Mo layer are deposited on three different glass substrates (size: 30 cm × 40 cm) with different Na concentrations: extra clear glass (solar grade opti-white diamond glass, Dia, high Na concentration), soda-lime glass (SLG, intermediate Na concentration), borosilicate glass (Boro, negligible Na concentration), see Fig. 1(a–c) 22,23,24. On Na-free borosilicate glass substrate, for comparative purposes, an additional 200 nm thick Mo:Na (3 wt%) film is deposited on top of the bilayer Mo stack, which serves as an extrinsic Nsource (see Fig. 1(d)). In addition, TiN and SiN thin films are utilised as sodium diffusion barrier and adhesion enhancement layer for the LP-deposited 500-nm Mo film onto SLG substrates (see Fig. 1(e,g)) 19. Finally, a 200 nm thick Mo:Na layer (see Fig. 1.1(f,h)), used as extrinsic Nsource, is deposited onto these rear contacts with a thin TiN or SiN sodium diffusion barrier layer.

The Mo films using different back contact designs were cut into 10 cm × 10 cm substrates for CuGa/In precursor deposition. A four-layer CuGa/In/CuGa/In stack design was used to prepare the Cu-In-Ga precursors by direct current (DC) magnetron sputtering. The selenisation of the precursors was performed in nitrogen atmosphere at ambient pressure condition. The heater temperatures of the RTP (Rapid thermal process) were 580 °C and the duration of the high temperature process was about 12 minutes. Thermally evaporated (420 °C) Se vapour was used as the Se source to form CIGS absorber layers. The average GGI ([Ga]/([Ga] [In])) and CGI ([Cu]/([Ga] [In])) ratios of the CIGS absorber measured by X-ray Fluorescence (XRF) were about 0.30 and0.93, respectively. Solar cells with these CGI and GGI ratios are reported to achieve high efficiencies of above 20% 25,26,27.

Results

In this study, CIGS absorbers with different Se contents (50 and 57 at.%) were fabricated via changing the Se crucible temperature in order to investigate the effect of the Se content on the formation of the MoSe2 layer. Figure 2 shows Scanning electron microscopy with energy dispersive X-ray spectroscopy (SEM-EDX) images of Se elemental mapping (at.%) across the absorber layer thickness. It can clearly be seen that the Se content in Fig. 2(a) decreases gradually from the top to the bottom of the absorber. At the rear surface of the absorber, i.e. at the Mo/CIGS interface, an approximately 200 nm thick CIGS layer was found to contain about 33% to 42% of Se, indicating a Se-poor absorber. However, for the sample shown in Fig. 2(b), the Se is distributed uniformly at above 50 at.%, showing a Se-rich absorber. The effects of the different rear contacts on the properties of the Se-poor and Se-rich CIGS absorber regions are investigated and discussed in the following sections.

Surface morphology

The surface morphology of the CIGS films formed on different rear contact structures was studied by SEM, see Fig. 3. The surface grain boundaries of Boro/Mo/CIGS and SLG/SiN/Mo/CIGS samples are clearly observed in Fig. 3(c,g), while white Na2Se precipitates are found on the surface of SLG/TiN/Mo/CIGS samples and samples having a Mo:Na layer at the rear contact as well.

The grain size of the CIGS films formed on bilayer Mo-coated diamond glass and soda-lime glass was much larger than that of the CIGS film formed on borosilicate glass, which was attributed to Na diffusion from the glass substrate through the bilayer Mo rear contact into the absorber. Thus, Na can diffuse through the voids between the columnar structures of bilayer Mo films. The addition of a 200 nm thick Mo:Na layer on top of the bilayer Mo rear contact slightly increased the grain size of the CIGS film. For the modified rear contacts deposited on SLG substrates, the grain size of the CIGS films formed on SLG/TiN/Mo rear contacts was much larger than that of CIGS films formed on SLG/SiN/Mo rear contacts (see Fig. 3(g)). In addition, the grain sizes of the Mo films deposited on SLG/SiN/Mo and on Boro/Mo substrates were quite similar (in the 0.5–1 μm range). These two findings indicate that a SiN barrier layer is more efficient in terms of blocking the Na diffusion from the soda-lime glass substrate than a TiN barrier layer. Using a Mo:Na film as extrinsic Nsource, the grain size of the CIGS films formed on SLG substrates with a TiN or SiN barrier layer were in the same range (about 1–3 μm) as those without a barrier layer in the rear contacts.

The grain sizes of the CIGS absorber were measured via surface SEM images. Histogram plots of the CIGS absorber grain size formed on different glass substrates and rear contacts are depicted in Fig. 4. The calculated average grain sizes are also shown in Fig. 4. The SEM studies clearly reveal that the Na incorporation either from the glass substrate or from the Mo:Na layer can enhance the CIGS grain size. It should also be noted that the CIGS film deposited on the SLG/TiN/Mo substrate showed the best lateral uniformity.

Chemical composition

The XRF measurement results for the CIGS absorber with 50 at.% Se formed on different rear contacts are shown in Table 1. The SeC and SeGI are [Se]/[Cu] and [Se]/([Ga] [In]) ratios. The average CGI and GGI ratios of these CIGS absorbers are about 0.93 and 0.30. The calculated Se/M ([Se]/([Cu] [In] [Ga]) ratio was slightly higher than 2.0. It is reported in the literature that the excess Se in CIGS solar cells exists in the form of a MoSe2 layer 20. The thickness of the MoSe2 ( \(\) ) layer can be estimated from the following equation 20 :

where \(_^\) is the atomic percentage of Se in the absorber as measured by XRF, \(_^\) is the nominal atomic percentage of Se in the absorber, \(\,_^_\) is the atomic percentage of Se in the MoSe2 (i.e. 0.67), \(_\) is the atomic density of CIGS (i.e. 4.23 × 10 22 cm −3 ), and \(\) is the atomic density of MoSe2 (i.e. 5.15 × 10 22 cm −3 ). The nominal atomic percentage of Se in the absorber \(_^\) is calculated by:

where x is the GGI ratio ([Ga]/([Ga] [In]) in the absorber.

Table 4 XRD measurement results of Se-rich CIGS absorbers formed on different glass substrates and different back contacts.

From Fig. 8(e), it can be seen that for the samples formed on borosilicate glass substrates with an additional Mo:Na layer as extrinsic Nsource, no MoSe2 diffraction signal was detected from the XRD measurement, which is due to the low amount of MoSe2 at the Mo/CIGS interface. The low growth rate of the MoSe2 layer can be explained by insufficient Na2Se intermediate compound formation at the rear surface. As discussed in our previous publication 18. for the CIGS films with higher Na diffusion from the substrate, it is observed to have higher Na and Se content in the form of Na2Sex compound at the Mo/CIGS surface 14,18,30. Additionally, from XRF measurements (Table 4), the observed MoSe2 thickness was significantly higher for Na-rich glass substrates than for Na-deficient ones. Thus, it is reasonable to conclude that the growth of the MoSe2 layer can be accelerated by Na2Sex. In other words, Na2Sex can be considered as a catalyst to form a thick MoSe2 layer.

If all the Na from the 200 nm thick Mo:Na capping layer is assumed to be uniformly distributed in the completed CIGS solar cell, the Na content in the CIGS absorber would be about 0.5 at.%, which is higher than the observed Na content in CIGS solar cells formed on SLG/Mo substrates 18. Thus, it can be concluded that most of the Na atoms diffusing from the Mo:Na layer are trapped within the Mo grains and cannot easily diffuse into the CIGS film.

From the XRD measurements, the value of the MoSe2 peak intensity ratio I(100)/I(110) was calculated, see Table 4. Since this value is higher than 1.0, this indicates a preferred grain orientation of the MoSe2 layer in the direction. It has been reported that a MoSe2 grain orientation in the direction reduces the contact resistance between the Mo and the CIGS layers, and enhances the adhesion of the CIGS layer to the rear contact 13,31,32,33. In addition, the conductivity of (100)-orientated MoSe2 grains, aligned in parallel to the Mo rear contact surface, is two times lower than that of (110)-orientated MoSe2 films, which have grains that are perpendicular to the Mo rear contact. Therefore, the MoSe2 intermediate layer formation needs to be adjusted by optimizing the selenisation process and by controlling the Na diffusion to obtain the preferred (110) orientation.

Discussion and Conclusion

Dependence of morphology and microstructure of the CIGS absorber on the Mo-coated glass substrates

The sodium diffusion either from the glass substrate or from the Mo rear contact plays a large role in the grain growth of the CIGS absorber layer, as the sodium can promote the grain growth of the CIGS absorber layer. Na diffusion to the front surface of the CIGS absorber is beneficial for forming large crystallites near the front surface. The CuInSe2 on the CIGS surface can be formed through the reaction of either CuSe2 or CuSe with InSe. However, the grain size of Cu2Se is larger than that of CuSe 34,35. The sodium selenide formed at the front surface of the CIGS can act as a Se source to promote the formation of larger CuSe2 grains, resulting in large CIS grains near the surface. At the same time, the diffused Na suppresses the CuInSe2 formation via reaction of smaller grain CuSe with InSe. The CIGS absorber formed on the TiN modified Mo rear contact shows the best uniformity in grain size, indicating the good control of Na diffusion from the SLG glass substrate by utilizing the TiN diffusion barrier layer.

All the CIGS films show a bilayer microstructure verified by XRD measurement, i.e., a Ga-rich CGS layer near the back contact and an In-rich CIS layer near the CdS buffer. The growth of the CuInSe2 and CuGaSe2 grains across the film thickness requires different formation energy, time and temperature during the selenisation process. The final Cu(In1−xGax)Se2 chalcogenide film is formed via inter-diffusion between the CGS and CIS layers. However, the Na diffusion from the substrate affects the inter-diffusion between the CGSand CIG layers. The CIGS absorbers deposited on borosilicate glass substrates show the highest x values, which are closer to the GGI ratio in the films measured by XRF. In contrast, CIGS films formed directly on the SLG substrate without any barrier layer show the lowest x value. The Na diffusion from the SLG substrate into the absorber reduces the In and Ga inter-diffusion, and thus reduces the Ga content at the front surface of the CIGS absorber. The TiN and SiN barrier layers can hinder the Na diffusion 18. and thus promoted Ga diffusion to the front surface. The lower x values of the samples deposited on Na-containing glass substrates indicate a lower Ga content at the front surface of the CIGS absorber.

Formation of MoSe2 intermediate layer

The precise control of the MoSe2 thickness and its crystal orientation are the most important factors determining the solar cell efficiency. Based on this study, the formation of the MoSe2 layer depends on both Na diffusion from the substrate and the Se content, which are related to the sodium content on the substrate, back contact structure and the selenisation process.

All selenisation processes in this study were conducted at a substrate temperature of 580 °C for a duration of 12 minutes. From the experiments, it was observed that there is no difference in the thickness of the MoSe2 layer in the Se-poor CIGS absorbers, while the thickness of the MoSe2 layer varies significantly in the Se-rich CIGS absorbers. In particular, for the Se-poor absorber, the Se is fully consumed by the metallic CIG precursor to form the CIGS absorber. Thus, the formation of MoSe2 is inhibited due to the lack of Se in the absorber forming environment. In this situation, the CIGS/Mo interface is stable even during long-time annealing in a Se-deficient condition, regardless of the amount of Na diffusing from the substrate.

In contrast, in Se-rich absorbers, Na2Se forming at both the Mo/CIGS interface and the CIGS grain boundaries tend to increase the MoSe2 formation probability. In this case, the excessive Se reacts with the Na diffusing from the substrate to form Na2Sex, which is usually located at the CIGS/Mo interface as a Se source and catalyses the MoSe2 formation. When there is insufficient or no Na available to form Na2Sex, it is less probable to form MoSe2 by the direct reaction between Se and Mo atoms. Hence, a very thin MoSe2 layer is found in the Se-rich samples made on rear contacts with low Na concentrations, for example the Na-free substrates and SLG substrates with SiN acting as a good Na diffusion barrier layer. A thick MoSe2 layer is only observed in samples with high Na diffusion from the substrate, and its thickness increases with increasing Na content in the glass substrate. However, if the MoSe2 layer gets too thick, the series resistance tends to increase and the solar cell efficiency was found to decrease.

In this study, the influence of the sodium diffusion from the substrate on the properties of two-step process formed CIGS absorber using elemental Se vapour are investigated. Using a TiN layer on SLG glass as a barrier layer to adjust Na diffusion, the CIGS film shows the best uniformity in grain size, indicating the good control of Na diffusion from the SLG glass substrate. The TiN barrier layers can partially reduce the Na diffusion, and thus enhance Ga and In inter-diffusion to form large CIGS grains. We also find that both Na and Se play crucial roles in the formation of the MoSe2 intermediate layer. A Se-rich atmosphere is required to grow a thick MoSe2 layer, while a Se-poor CIGS formation process leads to a thin MoSe2 layer, regardless of the Na content in the substrates. Since the series resistance of the CIGS solar cells increases with increasing MoSe2 layer thickness, an intermediate MoSe2 layer thickness will give the best device performance. Thus, the Se flow in the two-step CIGS formation process must be sufficiently high to obtain high-efficiency CIGS solar cells.

Acknowledgements

The Solar Energy Research Institute of Singapore (SERIS) is supported by the National University of Singapore (NUS) and the National Research Foundation (NRF) of Singapore through the Singapore Economic Development Board. This work was supported by NRF grant NRF2011EWT-CERP001-019, “CIGS solar cells and mini-modules: Establishment of an RD pilot line in Singapore and associated industry-relevant device research”.

Author information

Authors and Affiliations

• Solar Energy Research Institute of Singapore, National University of Singapore, Singapore, 117574, Singapore Weimin Li, Xia Yan, Armin G. Aberle Selvaraj Venkataraj
• Department of Electrical and Computer Engineering, National University of Singapore, Singapore, 117576, Singapore Weimin Li Armin G. Aberle
• Center for Information Photonics and Energy Materials, Shenzhen Institutes of Advanced Technology, Chinese Academy of Sciences, Shenzhen, 518055, China Weimin Li

Review on Substrate and Molybdenum Back Contact in CIGS Thin Film Solar Cell

Copper Indium Gallium Selenide- (CIGS-) based solar cells have become one of the most promising candidates among the thin film technologies for solar power generation. The current record efficiency of CIGS has reached 22.6% which is comparable to the crystalline silicon- (c-Si-) based solar cells. However, material properties and efficiency on small area devices are crucial aspects to be considered before manufacturing into large scale. The process for each layer of the CIGS solar cells, including the type of substrate used and deposition condition for the molybdenum back contact, will give a direct impact to the efficiency of the fabricated device. In this paper, brief introduction on the production, efficiency, etc. of a-Si, CdTe, and CIGS thin film solar cells and c-Si solar cells are first reviewed, followed by the recent progress of substrates. Different deposition techniques’ influence on the properties of molybdenum back contact for CIGS are discussed. Then, the formation and thickness influence factors of the interfacial MoSe2 layer are reviewed; its role in forming ohmic contact, possible detrimental effects, and characterization of the barrier layers are specified. Scale-up challenges/issues of CIGS module production are also presented to give an insight into commercializing CIGS solar cells.

Introduction

Copper Indium Selenide (CuInSe2 or CIS) is a ternary compound p-type absorber material belonging to the I-III-VI2 family [1]. The very first CIS material being synthesized was in 1953, and then, an efficiency of 12% has been reported for single crystal CuInSe2-based solar cells [2]. In 1976, the first CIS thin film solar cell with buffer layer CdS was fabricated with an efficiency of 4–5% by evaporating CuInSe2 powder in the presence of excess Se vapor (coevaporation) [3]. CIG-based thin film solar cell started to receive even more attention in 1981 when Mickelsen and Chen achieved an efficiency of 9.4% by using coevaporation technique from elemental sources [4]. From that onwards, numbers of emerged technologies such as alloying CIS with gallium (Ga) to become Copper Indium Gallium Selenide (CIGS), incorporating sodium (Na) into the CIGS absorber layer, and replacing thick cadmium sulfide (CdS) buffer layer with thin CdS layer have boosted the efficiency significantly. Copper Indium Gallium Selenide (CIGS) thin film solar cell currently holds a record efficiency of 22.6% since 2016 [5]. To accomplish the record efficiency, research institute ZSW has developed a new method to deposit a layer of potassium fluoride (KF) between CIGS and buffer layer (CdS) to improve the electrical properties of the solar cells, in particular the Voc.

The cell structure of CIGS is known as substrate configuration where the light enters the cell through Transparent Conducting Oxide (TCO), passes through the buffer layer, is absorbed by the CIGS, and then reaches the back contact, usually molybdenum, which is deposited on the substrate. The typical structure of CIGS solar cell consisting substrate/Mo/p-type CIGS/n-type CdS/i-ZnO/ZnO:Al/ARC/metal-grid is shown in Figure 1 [1]. The reason CIGS has been one of the most promising absorber layer for thin film photovoltaic devices is due to its high absorption coefficient for solar radiation and compatibility of its bandgap (1.6 eV–1.0 eV) [6]. The advantages of CIGS-based solar cells over CIS-based solar cells are as follows: (i) the bandgap can be tuned by adjusting the Ga/In ratio to match the solar spectrum. If all indium (In) is replaced by gallium (Ga), the CIGS bandgap increases from about 1.04 eV to 1.68 eV [7]. It has been stated that CIGS absorber layer can absorb most parts of the solar spectrum with a thickness of 1 μm [1]. Hence, a layer thickness of ~2.0–2.5 μm will be sufficient for the completed device, and a thinner layer device means reduction in raw material usage and lower production cost incurred. (ii) Ga incorporation can also improve the open-circuit voltage Voc of CIGS since Voc~Eg/2 (Eg is referring to bandgap) [1].

over, CIGS thin film solar cell has very high potential to overcome the cost level of conventional PV crystalline silicon (c-Si) technology [8]. The c-Si modules with efficiencies of 19–23% will have a production cost of 0.6–0.7/Wp [9]. Whereas for CIGS modules, manufacturing cost of 0.75/Wp can be achieved at 50 MW/yr production capacity with an average efficiency of 12% [10].

The substrate in CIGS has a crucial role in the development of the whole device. Deposition of the molybdenum back contact on rigid or flexible substrate will define the selenization condition. It has to be considered that flexible substrate cannot withstand high process temperature over 500°C, but high process temperature is required in crystallizing the CIGS absorber layer. Lower temperature processes have to be developed when dealing with flexible materials. Recently, an efficiency as high as 20.4% was achieved on flexible polymer substrate [11]. This has revealed that the flexible solar cells with performance close to rigid solar cells can be developed.

The following layer in CIGS after substrate is the molybdenum (Mo) back contact which acts as an optical reflector to reflect the light back to the absorber layer in CIGS solar cell [12, 13]. Molybdenum (Mo) is a preferred back contact material for CIGS solar cells because it does not react strongly with CIGS; it forms low-resistivity ohmic contact to CIGS, and the conductivity of Mo does not degrade during deposition of CIGS at high substrate temperature [14–17]. Mo has high conductivity and is more chemically stable and mechanically stable during CIGS growth (selenization) than other materials such as W, Ta, Nb, Cr, V, Ti, and Mn [18–23].

The layer after Mo back contact is Copper Indium Gallium (CIG) before going through the process of selenization. During the selenization process, selenium (Se) vapor will react with CIG to become CIGS and react with Mo to form the MoSe2 layer. This interfacial layer between Mo and CIGS is beneficial in terms of having a wider bandgap (1.35–1.41 eV) than CIGS, hence it can absorb more near-infrared light to improve the cell performance [24]. The formation of MoSe2 layer does not depend only on the selenization condition but also on the properties of the Mo film [25, 26]. Therefore, improving the properties of the Mo film can promote the growth of MoSe2 layer. Recently, substrates used in CIGS, either rigid or flexible, together with the properties of the Mo back contact and the MoSe2 interface layer were discussed in various papers [8, 25, 27, 28]. This paper aims to FOCUS on the mentioned area by first providing an overview of comparison between conventional PV and thin film solar cells followed by reporting the recent progress of the substrates in CIGS, specifically regarding the available substrates. This paper will then converge towards the Mo layer in CIGS and further discuss about the deposition techniques and effect of deposition condition on the properties of Mo back contact. Then, the formation and thickness influence factors of the interfacial MoSe2 layer will be reviewed in this paper. Scale-up issues of CIGS module production will also be presented to give an insight into commercializing CIGS solar cells.

Brief Introduction on the Production, Efficiency, etc. of a-Si, CdTe, and CIGS Thin Film Solar Cells and c-Si Solar Cells

Crystalline silicon (c-Si) currently plays a major part in thin film energy production with the highest module efficiency. Thick and rigid c-Si wafer (180 μm) is required for a good absorber (high absorption) in module production [1]. This is because silicon (Si) is a poor absorber due to its indirect bandgap nature and its low absorption coefficient (10 4 cm −1 ) [40]. Whereas for CIGS, its absorption coefficient is beyond 10 5 cm −1 making the thickness of CIGS to be 100 times less than the thickness of c-Si wafer. In addition, the module production of CIGS solar cell requires a lower thermal budget (~550°C) than the c-Si solar cell (~1100°C) [1, 41]. CIGS solar cells provide an alternative to Si solar cells, and it is highly competitive as less raw material, time, and cost are involved in module production. On the basis of material, less material usage will lead to lower manufacturing cost for CIGS solar cell and hence inducing a shorter energy payback time (~1 year) as compared to the c-Si solar cell (~2 years) [10, 42]. Aside from the conventional c-Si cells, there are also other commercially available thin film modules in the market, including amorphous silicon (a-Si), Cadmium Telluride (CdTe), and, as mentioned, Copper Indium Gallium Selenide (CIGS).

Amorphous silicon (a-Si) has a direct bandgap and only uses 1% of the material (Si) needed for crystalline silicon cells production [43]. a-Si module can be made flexible and lightweight which later enable various possibilities when it comes to application such as mounting on uneven surface, incorporating into small devices, and being portable. One of the a-Si solar cell advantages is the high actual power output in hot climate by having a low temperature coefficient [44, 45]. However, a-Si solar cells have difficulty in the solar market because the price of conventional PV (c-Si) has been decreasing dramatically [46]. The main issue in a-Si solar cell technology is having a low conversion efficiency. The 14% stabilized research efficiency (by National Institute of Advanced Industrial Science and Technology, AIST) and 12.3% module efficiency have clearly shown its limitation to compete in the PV sector. Until now, a-Si technology is matured and commonly being used in the application of consumer products (e.g., calculators, watches, and other noncritical outdoor applications) [43].

Cadmium Telluride (CdTe) currently holds the highest module efficiency at 18.6% due to one of its advantages of nearly ideal bandgap (1.45 eV) for solar terrestrial photo-conversion [36, 47]. The module production status of CdTe is currently ahead of a-Si and CIGS thin film solar cell as it can be produced at a cost of 0.75/Wp (watt peak), and the production cost at year 2020 is expected to be around €0.5/Wp [48, 49]. Amorphous silicon (a-Si) module can be produced at €0.75/Wp but in a lower efficiency (~9%) while CIGS is still trying to lower its production cost to below €1/Wp [49]. Cadmium Telluride (CdTe) is a favourable technology in terms of prospects in PV sector except the public has raised concern about the core material used in CdTe, cadmium, which is extremely toxic. It has been stated that cadmium forms a very stable compound with tellurium and hence is not banned as a hazardous substance, but before forming a compound, cadmium itself can lead to a variety of adverse health effects including cancer [50]. advance technology is needed to handle the cadmium and hence creating another challenge for CdTe in lowering the production cost.

Copper Indium Gallium Selenide (CIGS) solar cell is one of the best thin film candidate to look into because its lab scale efficiency (22.6%) has just surpassed CdTe’s lab scale efficiency (22.1%) based on the latest research cell record efficiency chart reported by NREL [37]. For CIGS solar cells, a large efficiency gap occurred between the lab scale cells (22.6%) and commercially available modules (15.7%). Although the CIGS module production presently lags behind of CdTe module due to process complexity, nonetheless, CIGS technologies still have a higher efficiency potential versus CdTe, and this can potentially make them more cost-effective than CdTe solar cells [1, 48]. The theoretical efficiency limit for CIGS bandgap of 1.14 eV is 33.5% [51]. According to the CIGS current record efficiency (22.6%), cell efficiency as high as 25% can be reached in the near future. In terms of module efficiency, CIGS (15.7%) is definitely not far behind from CdTe (18.6%) (Table 1).

Brief introduction on the production, efficiency, etc. of conventional PV (c-Si) and thin film solar cells (a-Si, CdTe, and CIGS) [10, 32, 33, 36, 37, 40, 42, 51, 52].

Recent Progress of Substrates

Soda-lime glass (SLG) is a type of rigid substrate being used widely in the CIGS thin film industry due to its material properties which can supply sufficient amount of Na to the absorber during coevaporation or selenization process [53–56]. The supply of sodium (Na) at 0.1 at% (atomic percentage) is reported to be beneficial for CIGS solar cells in terms of increasing the open-circuit voltage and fill factor that lead to an enhancement in solar cell efficiency [53, 57]. The improvement in device efficiency is mainly due to sodium (Na) that passivates the defects at the cadmium sulfide (CdS) and CIGS junction [58–60]. In addition, SLG meets most of the requirements needed such as good adhesion, low weight, and able to work on suitable temperature. SLG substrate also has an optimal coefficient of thermal expansion (CTE) for CIGS which is desired between 5 × 10 −6 and 12 × 10 −6 K −1 to avoid adhesion problem or crack formation during deposition of CIGS at high temperature [27]. Solar cell company like Solibro has achieved a current record efficiency of 21% (single-junction terrestrial cell) and 18.7% on minimodule using SLG as a substrate [61, 62]. The first time ever, in 2010, the substrate involved in pushing the efficiency of CIGS beyond 20% was SLG [63]. Other than the standard glass SLG, specialty glasses have been explored by several research groups which mainly FOCUS on high-temperature glasses to avoid softening of the substrate (SLG) during CIGS absorber deposition [64, 65]. High temperature condition will not only enhance the growth of absorber layer but also favour alkali-diffusion from substrate to absorbers. Thus, specialty glass serves as a medium to provide fine control of alkali-diffusion without softening at high temperature [66, 67]. The CIGS world record efficiency back in 2014, 21.7%, and current world record efficiency since 2016, 22.6%, were both achieved by the same research group ZSW on a specialty alkali-aluminosilicate glass that incorporate Na during the CIGS growth process [5, 61, 68].

In CIGS solar cells, the most commonly used flexible substrates are metals and polymers [27, 60]. The essential aspects to be considered for a suitable flexible substrate for CIGS films are dependent on different physical and chemical properties, such as thermal stability, vacuum compatibility, suitable coefficient of thermal expansion (CTE), humidity barrier function, chemical inertness, and surface smoothness [57]. Metals are able to withstand very high deposition temperature but they possess a rather high roughness, density, and CTE especially Al and Cu [27]. Furthermore, most of the metals like steel contains metallic impurities (Fe and Cr) that affect the device performance [27, 69, 70]. Therefore, a metal oxide barrier is used in order to provide electrical insulation between the substrate and the Mo back contact at the same time serves as a diffusion barrier against impurities from the metal substrate [27, 69, 71, 72]. Typical examples of barrier layer materials are Al2O3 and SiOx [73–75]. It has been announced that the highest cell efficiencies reached so far using stainless steel (SS) and titanium (Ti) foil as substrate are 17.7% and 17.9% [71, 76, 77]. Whereas the maximum cell efficiency (area of 0.5cm 2 with antireflective coating) achieved so far on an enamelled steel is 18.6% by ZSW research group [78]. These results prove the potential of metal as an alternative substrate to rigid glass.

Polymers as substrates have a much lower density, roughness than metals [27]. It allows direct monolithic integration of solar cells and roll-to-roll deposition process that can reduce manufacturing cost [27, 60]. Other than that, polymer has a high power-to-weight ratio and excellent radiation-hardness which makes polymer an ideal candidate for space application, but polymers cannot sustain high temperature of 550–660°C due to their limited thermal stability [79, 80]. Thus, low process temperature is required, and this generally leads to deterioration of absorber quality [81, 82]. Polyimide films are one of the few polymer films that can sustain temperature close to or above 450°C but not more than 500°C for a short period of time, but this type of polymer has high CTE and it varies depending on suppliers [27, 60]. Despite the challenges faced in using polymer foils as a substrate, the Swiss Federal Laboratories for Material Science and Technology (Empa) has successfully developed thin film solar cells, CIGS, on flexible polyimide foil with an efficiency of 20.4% [83]. The Empa research group was able to modify the properties of the CIGS layer so that it can be grown at low temperature without compensating the light absorption of the CIGS layer which contributes to the photo-current in solar cells [11]. Since 2010 until 2016, ceramics have also been used as flexible substrate [75, 84]. A submodule efficiency of 15.9% has been achieved by AIST in Japan using flexible zirconia-based ceramic sheet as a substrate [75]. Ceramic substrate is able to withstand higher temperature than the soda-lime glass and polymers due to its higher chemical, mechanical, and thermal stability and low porosity [84, 85]. However, its brittle behaviour might be an issue for industrial production on large scale (Table 2).

Summary of current record efficiency on different substrates (CIGS solar cell) [5, 11, 61, 62, 71, 75, 76, 78].

Different Deposition Techniques’ Influence on the Properties of Molybdenum Back Contact for CIGS

Current deposition techniques available in the thin film industry for Mo are Ion-beam sputtering, Direct Current (DC) sputtering, Radio Frequency (RF) sputtering, and also High-target-utilization sputtering (HiTUS) [25, 86–88]. The common techniques being used are DC and RF sputtering as they have been discussed profoundly in several papers [12, 86]. Since the characteristics of Mo thin film depend strongly on deposition method and deposition parameters, therefore, a comparison has been made to investigate the properties of Mo layers based on DC and RF sputtering techniques. DC sputtering is a well-established industrial process with high throughput capability and requires a cheaper set up cost, whereas RF sputtering deposits a more reflective Mo thin films but with added expense on RF power supplies and impedance matching networks [12, 14, 18, 88, 89]. DC sputtering technique has higher deposition rate than the RF sputtering technique, and DC sputtered films possess good uniformity and adhesion properties over large surface areas [24]. Low deposition rate that arises in RF sputtering technique is due to the number of cycles involved during deposition where RF sputtering only deposits in the second cycle of the AC supply. However, RF sputtered films are found to be more conductive with improvements in open-circuit voltage (Voc) and short circuit current (Isc) [90].

A desirable Mo back contact for CIGS solar cells is addressed to conductive, stress-free, well-adherent, uniform, and crystalline molybdenum (Mo) thin films with preferred orientation (110) on large area glass substrates [91]. By altering the deposition condition on discharge power, working pressure, substrate temperature, and target-to-substrate distance, the desired properties of Mo back contact in terms of physical, optical, and electrical can be obtained [20, 91, 92]. In this paper, the properties of the Mo back contact on both DC and RF sputtering will be further reviewed. The FOCUS deposition parameters will be on discharge power and working pressure, as these parameters have major effect on the properties of the DC and RF sputtered films.

In DC sputtering at low working pressure, surface with dense microstructures is observed while surface with loosened microstructures is observed at high working pressure [18, 55]. At low working pressure, atoms that obtained high kinetic energy are able to travel and bombard onto the substrate surface to form a compact Mo layer, whereas at high working pressure, insufficient kinetic energy has caused the number of atom bombardment towards the substrate surface to reduce, thus resulting in porous microstructure. The high-pressure Mo with loosened microstructure will eventually lead to greater inclusion of impurities such as oxygen that can increase the resistivity of the Mo film. The direct effect of working pressure of RF sputtering on Mo microstructure has not been reported yet, but low RF power was stated necessary to minimize stress and obtain a compact Mo microstructure despite high RF power deposits good quality of Mo [93]. The Mo deposited with high RF power caused delamination of the absorber layer after the selenization process, and this can be due to the presence of microstresses that existed on the Mo layer.

Another physical property to evaluate is the grain size of Mo. The trend of grain size is found to increase by increasing sputtering power and decreasing sputtering pressure in both DC and RF sputtering Mo films [88, 94–96]. In high power and low working pressure, the increase of Mo grain size causes space between grains to reduce and thus correlates well to the formation of densely packed Mo microstructure. Additionally, along all applied pressures, the grain size of DC sputtered Mo film is always larger than the RF sputtered Mo film [97]. One possible way to explain the formation of large grain size is probably higher power will induce higher flux, and high deposition rate of DC sputtering tends to increase the probability of the Mo particles to nucleate with each other.

Crystal structure of sputtered Mo back contact also plays a fundamental role in Mo quality determination. The lower the working gas pressure, the better the crystallinity of the DC sputtered film [88]. At lower working gas pressure (higher gas power), the atoms gain higher energy due to lesser scattering, and the atoms will then impact the substrate surface with sufficient energy which enhance the atoms mobility in order to facilitate atom diffusion and microvoid fill up, thus creating a conducive requirement for large grain growth and better crystallinity. The same crystallite behaviour was observed in RF sputtered films at low working pressure, but the degree of crystallization appeared to be lower as compared to the DC sputtered films [95, 97]. This can be related to the deposition rate of both RF and DC sputtering technique. Lower deposition rate of RF sputtering generally requires longer deposition time to achieve the same thickness as DC sputtered films, and within the deposition period, impurities such as oxygen can be introduced which can potentially restrain the process of atom diffusion and deteriorates the crystallite property of Mo film.

Surface morphology is one of the properties to be investigated, as it will affect the adhesion, optical, and electrical properties of the solar cell [20]. The surface roughness of DC and RF sputtered Mo films were found to be increased as the working pressure increased [24, 95]. At lower working pressure, the average roughness was due to large grain size formed under high kinetic energy gained by the atoms whereas the roughness at higher working pressure was attributed to loosened microstructure with voids formed under low energy excitation. Although both DC and RF sputtering portray similar surface roughness trend, RF sputtered films possess smoother surface morphology than the DC sputtered films at same working pressure [86]. Smoother morphology also indicates less void formation and porosity effect which contribute to a higher optical reflection in the Mo film; hence, RF sputtered films are more reflective than DC sputtered films as proven in several literatures [93, 98]. In other words, Mo films deposited under low working pressure will have higher optical reflection and conductivity due to densely packed microstructure and large grain size formed (lesser void formation yield to reduced incorporation of foreign atoms).

Furthermore, sputtering pressure has a major influence on the sheet resistance of DC and RF sputtered Mo films over sputtering power [12, 88]. As the working pressure increases, the sheet resistance increases and vice versa, whereas, the sheet resistivity of the Mo film is inversely proportional to the sputtering power [24]. The DC sputtered films also have higher resistivity as compared to the RF sputtered one in all applied working pressures [86]. This phenomenon is interrelated to the microstructure and surface morphology properties. As DC sputtered Mo films are rougher, it verifies the formation of larger voids (more porous microstructure) along the grain boundary, and this will allow impurities such as oxygen to occupy the voids resulting in an increase in resistivity. However, low-resistivity Mo films deposited under low working pressure often gets delaminated, and Mo deposited at high working pressure has better adhesion but less conductive [24]. Mo back contact that exhibits low resistivity and adhesive properties are challenging to simultaneously achieve in single Mo layer [20]. Thus, Mo multilayers like bilayer structure are usually applied to deposit a well-adhered and conductive Mo layer [99–101]. This was done by depositing the first bottom layer with high working pressure (to achieve better adhesion) and the second top layer at lower working pressure (to achieve lower resistivity).

Besides that, single Mo layer can also cause excessive tensile or compressive stress which lead to surface cracking and delamination between the Mo film and the substrate [18]. Bilayer structure is used as a solution to reduce excessive residual stress that exists on the Mo layer [24, 102]. The bottom Mo layer of tensile stress (formed under high pressure) and top Mo layer of compressive stress (formed under low pressure) contribute an overall residual stress compensation to avoid delamination of the Mo film from the substrate. It is an undeniable fact that depositing two layers of Mo can resolve adhesion and resistivity problem until certain extent, but taking into consideration of other properties such as microstructure, grain size, crystal structure, and surface morphology will further improve the quality (electrical and optical) of the Mo films. Referring to Table 3, RF sputtered Mo films are more reflective and conductive while DC sputtered Mo films have larger grain size, better crystallization, and better surface morphology. A combination mode of both sputtering techniques to form multilayer molybdenum (DC/RF) can be realised to optimise the Mo back contact accordingly. The potential of simultaneous DC and RF sputtering on Mo deposition can also be examined as this technique has been used by employing an induction coil in the DC path (to avoid short circuit to the RF voltage) for depositing ZnO/ZnO:Al window and contact layer with an improvement in Voc and Isc [90] (Table 3).

Formation and Thickness Influence Factors of MoSe2 Interface in CIGS Solar Cell

MoSe2 is an interface formed between molybdenum back contact and CIGS absorber layer during the process of selenization at high temperature (above 723 K) [103, 104]. The selenium diffuses into the Mo back contact and reacts to form MoSe2 along the process of selenization [105]. The MoSe2 layers consist of polycrystalline grains with columnar structure and lattice spacing which improve the adhesion between Mo and CIGS [106–108]. Instead of that, MoSe2 has a wider bandgap of 1.41 eV than the CIGS absorber that forms a back surface field which can hinder the recombination of electrons and holes [109–111]. The CIGS/Mo heterocontact including the MoSe2 layer leads to a favourable ohmic-type contact by the evaluation of dark I-V measurement at lower temperature [25, 109]. Without the interface layer, a Schottky contact will be formed at the Mo/CIGS contact, causing significant problem in resistive losses [112]. However, excessive formation of MoSe2 can lead to the delamination of the film and adverse effect on Voc and FF of the completed CIGS solar cells due to high resistance of the MoSe2 [113, 114]. Therefore, a range of specific MoSe2 thickness between 100 nm and 200 nm is required to ensure good adhesion and electrical contact between Mo/CIGS [115–117].

The thickness of the MoSe2 layer can be influenced by several factors such as sputtering conditions [117], residual stress in Mo layer [118], characterization of the barrier layer (TiN, MoNx, and MoOx) [115, 116, 119] and selenization conditions [120–122]. The quality of the Mo crystals improves with sputtering power, thereby reducing the resistance of the Mo back contact. As the sputtering power increases, the thickness of MoSe2 also increases [117]. This is due to the fact that increase in sputtering power enhances the diffraction intensity and crystallinity of Mo (110), (211) which facilitate the transformation of the cubic crystal structure Mo back contact into a hexagonal crystal structure MoSe2 layer [123, 124]. However, varying the sputtering power to achieve the desired MoSe2 thickness is not practical since high and low sputtering power are required to deposit a well-adhered and conductive Mo layer (bilayer structure) [99]. The MoSe2 layer thickness is also dependent on the in-grain density of Mo which is interrelated to the residual stress. In a typical sputtered Mo layer, the tensile residual stress of Mo increases with the sputtering pressure up to 10 mTorr and decreases with further increase of pressure [118]. The increase in pressure reduces the mean free path and decreases the ion energy of the Mo particles, causing an increase in tensile stress (Region I). As the pressure further increases beyond 10 mTorr (Region II), the atomic attraction across gain boundary increases, creating porosity which eliminates the grain boundary attraction, thus reducing tensile stress. The reduction of the tensile stress in Region II will increase the in-grain density of the Mo back contact and decrease the MoSe2 reactivity (thickness) during the selenization process.

In recent years, barrier layers are used to control the excessive formation of MoSe2. A thin Titanium Nitride (TiN) barrier layer with a thickness of ~20 nm was grown on a Mo-coated soda-lime glass under vacuum condition prior to CZTSe deposition (CZTS has the same solar cell structure and works similarly to CIGS) [115]. Transition metal nitrides are used as a barrier layer because several researches showed that transition metal nitrides worked better in preventing diffusion at higher temperature than transition metals (Ta, Ti, W, and Mo) [125–127]. The TiN barrier layer has successfully suppressed the growth of MoSe2 from ~1300 nm to ~200 nm significantly improving the Voc, Jsc, and FF of the device. Other than TiN, the same concept has been applied by forming a molybdenum nitride (MoNx) barrier layer to passivate the Mo back contact against selenization [116]. With MoNx thickness of ~120 nm, the thickness of MoSe2 formed was ~150 nm without consuming the entire upper thin Mo layer (20–30 nm left). One previous study has also demonstrated that the MoNx barrier layer is able to control the formation of MoSe2 and improve FF and current-voltage characteristic of the device, but thicker barrier layer will increase the series resistance of the device [128]. Other than the transition metal nitride barrier layer, molybdenum oxide (MoOx) was proposed to control the growth of MoSe2 [119]. It has been reported that a thin layer of MoO2 as low as 10 nm is able to prevent the Mo back contact from overselenization and improve the Voc, FF, and shunt resistance (Rsh) of the device [129]. In addition, oxygen allows diffusion of alkali metals during selenization which will enhance the performance of the device [130–134]. The behaviour of the MoO2 barrier layer to control the formation of MoSe2 can be explained with their corresponding Gibbs free energy. At 900 K, the Gibbs free energy of the reaction between Mo and oxygen (−423 kJ/mol) is lower than the one of Mo with selenium (−129 kJ/mol) causing the reaction of MoO2 with selenium (294 kJ/mol) to be thermodynamically unfavoured, validating the passivation effect of MoO2 on Mo back contact against selenization [119].

A well-crystallised CIGS absorber layer requires a process temperature of at least 500°C and above for the selenization process [118]. Nonetheless, at high process temperature, the thickness of the MoSe2 layer will increase significantly [117]. If the process temperature is lowered to reduce the thickness of MoSe2, deterioration of the absorber layer (electrical properties) will occur [135]. Besides the mentioned thickness-influencing factors, sodium (Na) might be the dominant factor to control formation of MoSe2 layer [25, 123]. An experiment has been conducted by varying the amount of Na (using SiOx barrier layer) diffuses from the substrate across the Mo back contact to the CIGS absorber layer [123]. The experiment concluded that Na aids in the formation of MoSe2, but if the Na content is too high, it passivates the grain boundaries at CIGS layer to form Na2Sex; thus, lesser Se atom is available to react with Mo, and this will retard the formation of MoSe2. Comparing these thickness-influencing factors, barrier layer appeared to be the most effective method to control the thickness of MoSe2, allowing sodium (Na) diffusion while minimizing the adverse effect on cell performance.

Scale-up Challenges/Issues of CIGS Thin Film Modules

Since 2006, thin film solar cell production in the U.S. has outperformed the production of c-Si solar cell, becoming the least expensive technology to be manufactured [35, 48]. The existing thin film PV technologies especially CIGS have reached over 1.6 GW of cumulative module production in 2015 and is reported to have a high cost reduction potential at high production volumes [136, 137]. For a production capacity of 1000 MW/yr with 15% module efficiency, the CIGS module production cost as low as 0.34/Wp can be achieved [10]. Due to photovoltaics cost declining with maturity of the conventional PV technologies and new entry of China into the market in 2010, the future challenges for CIGS production will be combining high production volumes with high throughput, sufficient yield, and superior quality. Meaning that the production cost has to be brought down to make it competitive with conventional sources.

Currently, the leading CIGS module manufacturer Solar Frontier claims that their CIS module is able to output more electricity than the conventional crystalline silicon (c-Si) in real-world conditions such as better performance at high temperatures, low-light condition, light-soaking effect, and also shadow tolerance [138]. With mass production of modules having efficiencies ranging from 11.8% to 13.8%, Solar Frontier was able to achieve an annual production of ~1 GW and over 3 GW of shipments worldwide in 2015 [139, 140]. The statement made on surpassing the performance of c-Si solar cells is according to the cadmium- and lead-free CIS module (SF 150-170S Series) which provides an efficiency of 13.8% for a total area of 12,280 cm 2 [140, 141]. The current technology offered by Solar Frontier with high energy yield leads to a shorter energy payback time which fulfil the requirement to be competitive towards c-Si solar cells by producing high-quality module at lower cost. On the other hand, Solar Frontier has also attained a higher efficiency of 19.2% on cadmium-free CIGS minimodule [36, 142]. This research direction affirms the potential of cadmium-free CIGS module with consideration on environmental issues (reduce usage of cadmium) as the energy production increases.

The production of CdTe PV modules is ahead of CIGS PV modules at present stage due to simplicity of the process. However, different approaches have been used to boost up the CIGS cell efficiency such as applying Post Deposition Treatment (PDT) on the CIGS surface with alkali elements, incorporating more gallium into the absorber layer (CIGS) and also combining with other materials such as perovskites for multijunctions [1]. PDT is the process developed by ZSW to reach 22.6% efficiency in CIGS solar cells, but the compatibility and practicality between new technologies and scaling it up into mass production should be taken into account earnestly. Therefore, the challenges/issues in CIGS solar cell production are discussed as below: (i) Uniformity of CIGS absorber film over large areas: Uniformity is essential for electrical and optical properties of high-efficiency solar cells, and it directly influences the yield in production. The yield determines actual production volume in MW/yr and production cost in ¢/Wp [35]. Coevaporation can be an appropriate technique for large area substrate because the coevaporated CuGaIn precursor has a higher selenization rate than the cosputtered one, and it has been actively developed by many manufacturing companies [122] (ii) Standardizing the cell fabrication process: In the industry, Mo and ZnO films were sputtered (vacuum system); buffer layer CdS was deposited using Chemical Bath Deposition (nonvacuum system), and absorber layer CIGS was deposited with the coevaporation method (vacuum system) [143]. Fabricating the cell in and out from vacuum and nonvacuum process will cause difficulty in troubleshooting once a problem occurs. Substrate handling in open air environment and between different tools can contaminate the substrate. This can be avoided by using a vacuum process along the fabrication of the whole cell (iii) Presence of moisture in CIGS modules: Water vapor will oxidise the back contact molybdenum (Mo) causing Mo to degrade [8]. To solve this issue, a robust encapsulation technology with the properties of durability, adhesion, thermal stability, etc. is required (iv) Long processing time for CIGS and TCO layers: The processing time for CIGS (including deposition and selenization) and Transparent Conducting Oxide (TCO) layers should be reduced to meet the required industrial production time (10 min [1]) when it involves large volume production. The processing time can be improved by using a thinner CIGS absorber layer and high-speed deposition technique for TCO (v) Cadmium used in the CIGS buffer layer (CdS) is toxic, and disposal of the cadmium-containing product causes detrimental effect to human health [144]. over, the use of CdS buffer layer leads to optical absorption loss [145]. Hence, the CdS buffer layer can be replaced by other appropriate wider bandgap buffer materials to improve the short-circuit current (Jsc) of the device (vi) Indium scarcity on CIGS module: Production of indium currently relies in by-products of mining and refining of other material (in particular, zinc) [146]. There is a concern raised upon scarcity of indium might escalate the price and can be a threat to CIGS ambitions for production cost reduction and cost competitiveness in the wider PV market. Addressing this issue, CIGS layer of ≤1 μm should be used without compromising Jsc of the device [1].

The major manufacturer of CdTe thin films, First Solar, was able to obtain a LCOE at 0.15 /kWh, and the LCOE was targeted to be brought down to 0.08 /kWh with the system cost of 2 /W [154]. Whereas CIGS (14% efficiency) has attained a decent LCOE at 0.084 /kWh for a 100 MW scale PV system [155]. Besides that, the CIGS LCOE reported at 0.084 /kWh was being compared with c-Si solar cell (16% efficiency) in an identical location and system, and the LCOE of the c-Si solar cell was at 0.80 /kWh, lower than CIGS but comparable to each other [155]. This can be explained by the fact that c-Si solar cell has better module efficiency causing the Balance of System (BOS) to be reduced meanwhile CIGS offers better performance under high temperature and low-light condition thus diminishing the LCOE differences between c-Si and CIGS solar cell. The LCOE of CIGS is also believed to be higher than the CdTe thin films at this moment; it is because CIGS thin film possesses lower module efficiency and requires more complex system that tends to increase both cost on BOS and PV system. Nevertheless, CIGS technology is still feasible and promising as its record efficiency (22.6%) has surpassed CdTe (21.6%), and the theoretical efficiency limit of CIGS (33.5%) is higher than the c-Si (29.43%) solar cells [37, 51, 52]. This implies that CIGS solar cell with greater efficiency potential will result in lower material usage, lower system cost, and higher energy yield which eventually contribute to an exceptional Levelized Cost of Energy (LCOE).

Conclusion

CIGS solar cells are believed to have a very high potential against c-Si and CdTe solar cells in achieving low production cost with high module efficiency as the CIGS possesses better absorption coefficient (lower material usage), requires lower thermal budget than c-Si solar cells, and its record efficiency (22.6%) has just surpassed CdTe (22.1%). The rigid substrates such as soda-lime glass and alkali-aluminosilicate glass give rise to higher efficiency with direct Na incorporation. However, flexible substrates have proven its capability to be an alternative to rigid substrates by achieving a decent efficiency comparable to the rigid substrate solar cells by applying Post Deposition Treatment. DC sputtered Mo films are favouring physical properties while RF sputtered Mo films are favouring electrical properties. Thus, multilayer molybdenum (DC/RF) or simultaneous DC and RF sputtering mode can be explored to optimise the molybdenum back contact. On the other hand, the MoSe2 layer is important in forming ohmic-type contact and improving the adhesion between CIGS and Mo layers. Excessive formation of MoSe2 can cause delamination problem and increase in resistivity of the solar cell. Addressing this issue, the most effective method to control the thickness of MoSe2 is forming a barrier layer in between the bilayer structure Mo back contact. The scale-up challenges/issues of CIGS thin film modules discussed are possible to overcome as current CIGS technology has started to offer better performance at high temperature and low-light condition than the c-Si solar cells in real-world conditions. Lastly, with greater efficiency potential and higher total lifetime power produced, CIGS technology will eventually attain an outstanding Levelized Cost of Energy (LCOE) against c-Si and CdTe solar cells.

Conflicts of Interest

The authors declare that there is no conflict of interest regarding the publication of this paper.

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Copyright © 2018 Kam Hoe Ong et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Optoelectronics Simulation of CIGS-Based Solar Cells Using a Cd-Free Nontoxic ZrSxSe2−x as a Novel Buffer Layer

In this work, we investigate the performance of CIGS-based thin-film solar cells employing the SCAPS-1D simulation package. The paper is mainly devoted to the development of the ZrSxSe2−x (where 0 ≤ x ≤ 2) transition metal dichalcogenide (TMDC) as a Cd-free, nontoxic, and abundant buffer layer, the first of its kind. In the first step, we have evaluated the impact of the p-MoSe2 interfacial layer between the GIGS absorber and Mo back contact. The J–V characteristic showed a higher slope, revealing that the p-MoSe2 layer at the CIGS/Mo interfaces beneficially on the CIGS/Mo hetero-contact, mediating the quasi-ohmic contact rather than the Schottky type. For the optimized solar cell using the ZrSxSe2−x as a buffer layer, the photovoltaic parameters, such as the short-circuit current density, open-circuit voltage, Fill Factor, and efficiency, were investigated versus the thickness, carrier concentration, and bandgap values. The results reveal an optimum efficiency of ~ 25.5% at a bandgap of 1.3 eV, corresponding to ZrS0.8Se1.2 (i.e., x = 0.8) and 180 nm thicknesses, at a high carrier concentration of 1 × 10 18 cm −3. Furthermore, the solar cell performance is assessed with the increment of the operating temperature from 275 to 475 K. The observed decrease in the Voc is ascribed to the rise in the reverse saturation current associated with the higher temperatures. The study concludes an excellent potential for fabricating high-performance CIGS thin solar cells using a Cd-free nontoxic buffer layer.

Introduction

Chalcopyrite Cu(In,Ga)Se2 (CIGS) compounds have been drawing significant consideration as capable and appropriate candidates for cost-effective, efficient photovoltaic devices due to their remarkable physical and electrical properties [1,2,3]. They crystallize in a tetragonal chalcopyrite structure derived from the cubic sphalerite structure. One consequence of this structure is the possibility of varying its bandgap values (Eg). The Eg varies from pure CIS to pure CGS from approximately 1.0 to 1.7 eV, subject to gallium and indium percentages in the composition [4]. Additionally, they possess a high absorption coefficient in the range of 10 4 cm −1. which implies that 95% of the solar radiance with photons has energy above the Eg that can be absorbed within 1 μm of CIGS film, whereas, for example, crystalline Si might need more than 100-μm thickness to absorb that much of solar energy [5, 6]. High efficiency of 20.3% has been recorded for a laboratory-scale CIGS solar cell [7]. Other efficiencies of around 19.9% [8] and 17.4% [9] for the CIGS thin-lm sub-module have been achieved. over, in the CIGS-based thin-film solar cells, molybdenum is the most commonly employed material for the back contact. For an efficient CIGS solar cell, the Mo layer shall establish a low-resistance ohmic back contact with the CIGS layer. Else, substantial resistive losses due to the Schottky barrier at the CIGS/Mo contact occur. It has been found that forming a proper intermediate layer between the CIGS absorber and Mo provides low resistance quasi-ohmic contact. For example, a skinny p-MoSe2 layer might be grown on Mo back contact within the deposition of the CIGS layers, and the contact between Mo and CIGS reveals an ohmic character [10,11,12]. It is currently believed that for many solar cell structures, the existence of such an interfacial layer can promote a quasi-ohmic electrical contact across the CIGS–Mo interface. Accordingly, a p-MoSe2 intermediate layer is added to the investigated structure in the study.

Even though the CIGS has become an excellent absorber for thin-film solar cells, optimizing the buffer layer is still a dispute for improving cell performance. The buffer layer is an n-type material used to form the p–n junction of the photovoltaic devices. So far, the most common buffer layer in the CIGS solar cells is the CdS. Introducing a suitable buffer layer was a significant step in developing higher-efficiency solar cells [13]. However, the disadvantages of CdS cannot be disregarded. Cadmium used in the buffer layer is a toxic material with restricted use in electronics. Disposition of the Ca-containing outputs evokes injurious effects on human health, and appropriate treatment of large quantities of toxic waste might be an enormous and costly duty. Apart from toxicity, the bandgap of CdS is about 2.40 to 2.50 eV, which results in significant optical absorption fall in the short wavelength range, limiting the overall performance of solar cells. Because of the disadvantages of the CdS buffer, considerable efforts have been put into finding Cd-free alternative buffer candidates, e.g., ZnS, Ins, ZnSe, Zn(O, OH), Zn(O, S, OH), and ZrO [14,15,16,17]. The potentiality of other transition metal dichalcogenide materials (TMDCs), namely the MoS2 as a buffer layer for CZTS and CIGS solar cells, has been studied [17, 18]. This paper explores the potential of a novel material as a buffer layer, namely, the zirconium sulfoselenide (ZrSxSe2−x) TMDCs. They have a bunch of remarkable characteristics as 2D materials. Additionally, one can engineer their bandgap values between ∼ 1.2 and 2.7 eV, using what has been termed chemical architecture by changing the sulfur content, i.e., the x parameter. The TMDCs exhibit a broad spectrum of electronic properties varying from metals or semimetals through semiconducting to insulating [19, 20]. Also, the materials exhibit an unusual morphology as thin, flexible, high-quality surfaces with no dangling bonds, introducing their suitability for thin-film fabrications. The general expression of the TMDC material is MX2, composed of X–M–X planes, where a layer of metal atoms lies between two chalcogens layers with a strong covalent bonding for the in-plane sheets and a weak Van der Waals coupling held the sheets [21,22,23]. Furthermore, using environmentally friendly abundant, nontoxic, and cheap materials is essential for solar cell fabrications and valid for the TMDC semiconductors. over, the implementation of TMDC as a buffer layer resolves other parameters that might directly impact the cell performance, like the lattice parameter mismatch and the coefficient of the thermal expansion match between the absorber and the buffer layers. The advantage of TMDCs is that they have thermal and chemical stabilities of the n-type character associated with high carrier mobility.

In this study, our efforts have focused on a quantitative numerical study on the performance of a new proposed CIGS-based solar cell using the nontoxic n-ZrSxSe2−x as a buffer layer and a p-MoSe2 TMDC material as an interfacial layer between Mo and the absorber. The structure of the solar cell is Al/n-ZnO/i-ZnO/ZrSxSe2−x/CIGS/p-MoSe2/Mo/SLG. First, we touch upon the effect of the p-MoSe2 interfacial layer. The CIGS absorber thickness is optimized. Then, solar cell performance parameters such as open-circuit voltage Voc, short circuit current density Jsc, fill factor FF, and the conversion efficiency η are investigated as a function of the bandgap values, the thickness, and the acceptor carrier concentration (NA) of the ZrSxSe2−x buffer layer. Finally, the influence of the device operating temperature, in the range from 275 to 475 K, on the overall cell performance is investigated and discussed.

Device Modeling and Simulation

A numerical simulation is a powerful tool that leads to a comprehensive understanding of solar cells, identifying the main parameters that affect their performance and designing new experimental production structures. Many simulation programs have been used for the CIGS solar cells, such as SILVACO ATLAS [23], AMPS [24], and SCAPS [25, 26]. All simulations were obtained using the solar cell capacitance simulator (SCAPS) in this work. SCAPS is numerical simulation software that was prepared at the University of Gent [27]. It has been widely employed to simulate the electrical characteristics, performance, and spectral response of several crystallines and thin film–based photovoltaic cells, such as perovskite, CZTS, Si, CdTe, and CIGS solar cells. One of its advantages is that the SCAPS simulation results agreed well with the corresponding experimental results, such as perovskite, CZTS, Si, CdTe, and CIGS solar cells [26, 28,29,30,31]. SCAPS employs Poisson’s and continuity equations for electrons and holes to estimate the steady-state Band diagram, recombination profile, and carrier transport in one dimension. Figure 1a depicts the layer structure of the proposed model CIGS device. The Mo/MoSe2/CIGS/ZrSxSe2−x/ZnO/ZnO/Al structure is a heterojunction, i.e., a junction between semiconductors of different bandgaps. On the other hand, Fig. 1b shows the corresponding energy Band diagram of the proposed CIGS solar cell model using the ZrSxSe2−x buffer layer. The ZnO plays the role of the conventional window layer. The parameters used for the window layer are taken as the thickness of 150 nm, doping of 5 × 18 cm −3. and bandgap of 3.3 eV. For all simulations, this work applies the AM1.5 light spectrum illumination, and light intensity of 1000 (W/m 2 ) at a room temperature of 300 K. The surface recombination velocities are assumed to be \(^\) cm/s, and the coefficients of reflection of back and front surfaces are approximated to 10%. As per all simulation work, the input data for each material considered here have been collected and listed in Table 1 [32,33,34]. SCAPS simulates the energy diagrams in steady states and calculates the electron–hole recombination profile with the charge carrier transport in one dimension by solving the fundamental semiconductor equations. These equations are Poisson’s equations that connect the electrostatic potential to the charge (Eq. 1) and the continuity equations for electrons and holes (Eqs. 2 and 3):

Regarding the impact of the thickness of the p-MoSe2 interfacial layer, a correlated study is performed, considering the absorber layer thickness. The latter is one of the critical parameters that optimize solar cell performance. It should be carefully selected to absorb the highest number of photons and not too large to control and minimize the reverse saturation current. The obtained results of the photovoltaic parameters such as open-circuit voltage Voc, short circuit current density Jsc, fill factor FF, and the conversion efficiency η are shown in Fig. 3. The thicknesses of the absorber and the interfacial layers vary from 0.2 to 1.0 μm and from 0.05 to 0.2 μm, respectively. The results show an η of 23.48%, corresponding to an absorber thickness of 0.8 μm and an interfacial layer of thickness of 0.16 μm. A constant Jsc (i.e., horizontal curve) is observed with the increase of p-MoSe2 thickness. For example, its value is 33.32 mA/cm 2 at a thickness of 0.8 μm of the CIGS. This is because the Jsc parameter depends mainly on the change in the absorber thickness. On the other hand, Jsc increases from 31.69 to 38.16 mA/cm 2 as the CIGS absorber thickness increases from 0.2 to 0.85 μm. The parameter Voc has increased from 0.703 to 0.743 V, increasing the p-MoSe2 thickness from 0.05 to 0.2 μm, and the efficiency improves from 22.16 to 23.48%. Figure 3d reveals an optimum value of FF of 83%, which corresponds to maximum efficiency at the possible thinnest absorber thickness. Indeed, the results indicate an enhancement of the performance by increasing the thickness of the absorber layer. We observe an increase in the device efficiency of ~ 3.5% with the increase of CIGS thickness from 0.17 to 0.9 μm. Beyond this value of CIGS thicknesses, no dramatic improvement is observed. It can be concluded that a CIGS thickness of 0.9 μm with a high doping concentration of 1 × 10 18 cm −3 is enough to keep the device operating at its optimum efficiency since it is useless to consume extra materials. The above results are discussed regarding the back contact recombination and its relation to the absorber thickness. Many photogenerated carriers are likely to suffer from the recombination process at the back contact, and fewer ones can contribute to quantum efficiency. Increasing the absorber layer thickness leads to more photogenerated carriers than can be collected before the recombination process, and that is why the efficiency is enhanced. Additionally, the back contact recombination current density is reduced in a cell with a thick absorber layer since the back contact is placed farther from the depletion region. Nonetheless, we cannot increase the absorber thickness excessively, and optimization is still needed. Based on the simulation results, with a thickness of more than 0.9 μm, the efficiency only increases barely and is considered almost saturated. Indeed, with the absorber layer thickness increasing further, some incoming photons are absorbed deep into the absorber layer, away from the depletion region. The resulting carriers will not reach the space charge region during the assigned lifetime, recombining in the absorber bulk. In other words, increasing the thickness of the absorber layer raises the opportunity for bulk recombination. For the impact of the interfacial layer, a thinner p-MoS2 layer reduces the shunt resistance, and potential barriers would exist at the CZTS/p-MoS2 interface and p-MoS2/Mo interface. These barriers impede the drift of the photogenerated holes. Increasing the thickness of the MoS2 interfacial layer reveals an increase in the optical absorption, which in turn shows a higher photogenerated charge carrier. When the thickness of p-MoS2 is increased, the height of the barrier at the CIGS/p-MoS2 interface is reduced. This is because the work function and bandgap values of the p-MoS2 decrease with increasing thickness. Hence, the collection of photogenerated holes from CIGS to Mo is improved, providing better CZTS solar cell photovoltaic parameters. It is interesting to mention here that it has been noted that the MoS2 layer might form a back-surface field for the photogenerated electrons [39]. This back-surface field (BSF) causes a reduction in back-contact recombination and leads to improve cell performance, as it works as a hole transport layer (HTL) at the back contact. The relationship between the BSF and the recombination rate at the back contact has been investigated experimentally, e.g., [40].

Figure 4a–d depict the impact of the bandgap values of the ZrSxSe2−x buffer layer and its thickness in a contour plot in the range of 0.02 to 0.2 μm (y-axis) and from 1.3 to 1.9 eV (x-axis) on the photovoltaic performance parameters. The results show that by increasing the buffer layer thickness from 0.025 to 0.15 μm, one observes a neglected decrease in the Jsc from 37.85 to 37.80 mA/cm 2. This might be because the existing hole concentrations close to the semiconductor junction could increase the possibility of facing interface recombination [41,42,43]. It is worthy to say that bandgap values above 1.5 eV cause a detrimental effect on the cell performance. In buffer layers with a larger bandgap, photons cannot gain the energy needed to absorb the photons with higher wavelengths and, accordingly, will not generate electron–hole pairs. In other words, because of the buffer layer’s parasitic absorption, the absorber/buffer interface prevents the electrons from entering the junction [43]. A more quantitative picture of the results can be discussed by referring to the absorber-buffer interface. Considering the bandgap values for the CIGS/ZrSxSe2−x system, it is well accepted that a cliff-like Band alignment is assumed, where the valence Band maximum (VBM) and conduction Band minimum (CBM) of the ZrSxSe2−x buffer is less than that of the CIGS absorber. This develops the type-II heterojunction, a supportive structure for efficient charge separation [44]. Previous device simulations in literature have shown that the ideal conduction Band offset (CBO), in a moderate cliff-like, ranges from 0 to 0.4 eV [45]. When the value of the bandgap increases, most likely because of the downward shift in the valence Band Ev, this increase in valence Band offset (ΔEv) value blocks the charge separation promoted to the p-type absorber. Noting that, a slight increase in ∆Ec value will not dramatically affect the electron diffusion. Similar behavior was observed in previous studies on various chalcogenides of ZTSSe and CIGS [46, 47]. The optimum values obtained suggest an optimum efficiency of about 25.08% at a low bandgap energy of 1.3 eV, corresponding to buffer layer thicknesses of 180 nm. The optimum values of the FF, Jsc, and Voc were observed of 85.17%, 37.87 mA, and 0.77 V, respectively. According to our previous investigation of the bandgap values as a function of the sulfur content (i.e., the x parameter), we conclude that the obtained optimum bandgap corresponds to ZrS0.8Se1.2 [48]. Many studies recommended an optimal thickness of the CdS buffer layer to be between 50 and 100 nm [49], and the solar cell performance decreased out of this range. However, this limitation of the ZrS0.8Se1.2 buffer layer thickness is relaxed since we introduce a nontoxic, abundant, and Cd-free alternative.

Furthermore, the influence of buffer carrier concentration of ZrSxSe2−x on the solar cell performance parapets is studied. The doping level of the buffer is assumed to vary from 1 × 10 11 to 1 × 10 18 cm −3. The impact of increasing the carrier concentration on the photovoltaic parameters is demonstrated in Fig. 5. The results reveal a significant increase in the cell performance at doping levels above 1 × 10 16 cm −3. The Jsc is almost constant at about 37.875 mA/cm 2. Since the absorber thickness is held constant, no change in Jsc is observed. According to the literature, this parameter is most affected by the decrease in absorber thickness. The observation agrees with previous results using other TMDC buffer layers such as the MoS2 [17, 18]. The solar cell performance parameter FF increases by 11%, i.e., from 75.4 to 84.1%. At the same time, an increase in FF will collectively increase the conversion efficiency of the solar cells. The efficiency enhances from 22.88 to 25.52% with the increase of the doping level from 10 12 to 10 19 cm −3. It is recommended to have heavily doped with a donor concentration of 1 × 10 18 cm −3 to keep the device operating at optimal efficiency of 25.52%. The results can be discussed as follows. The highly doped layer produces the suited conduction Band downward bending at the interface and benefits the photoelectron collection [50]. As the carrier density of the buffer layer increases, more atoms can be ionized, and accordingly, more electrons are photogenerated. Besides, the electrostatic potential barrier in buffer/absorber ZrSxSe2−x/CIGS and window/buffer ZnO/ZrSxSe2−x heterojunctions decreases and allows an enlargement in the space charge zone, improving the collection of the photogenerated carriers, and increasing the conversion efficiency output can be achieved. The results are consistent with previously reported behavior, showing that the conversion efficiency increases with the carrier densities [18].

Figure 6 illustrates the obtained optical response of the ZrSxSe2−x buffer layer thickness. From Fig. 6, we can say that the transmittance of the sample appears to be nearly independent of the buffer layer thickness. However, an improvement of QE from 86.1 to 89.6% was observed at longer wavelengths ranging from 700 to 820 nm. The inset of Fig. 6 is the enlargement of the observed results in the investigated range. One acceptable explanation is that increasing the buffer layer thickness leads to the generation of more charge carriers. This wavelength matches the bandgap energy of the buffer layer. The enhancement behavior is due to the expected greater amount of atoms ionized to photogenerate more electrons.

Finally, this section focuses on the device operating temperature effect on the overall cell performance. The temperature was tuned in the range of 275 to 475 K. The simulated results are illustrated in Fig. 7. The device maintained a very high performance with device efficiency greater than 24% at operating temperatures below 340 K. A significant device performance drop in all performance parameters was observed at operating temperatures above this value. The Voc, FF, and efficiency values decrease with increasing the temperature while Jsc reduces slightly. Jsc decreases from 37.820 to 37.40 mA/cm 2 as the temperature increases from 340 to 475 K. The conversion efficiency decreases from 17.55 to 12.89%, about a 9% drop. The CIGS solar cell devices obtained J–V characteristic curves at different temperatures are shown in Fig. 7b. It displays a significant increase in the current density of the operating temperature in the range of 275 to 475 K. The observed decline in the Voc is due to the increase in the reverse saturation current associated with the temperature rise. The decrease in the Voc with the increase in temperature is validated by Eq. (4), which relates the Voc and the saturation current Io [51, 52]:

The increase in the saturation current density is mainly due to the increase of the intrinsic carrier concentration ni and the decrease in the bandgap of the semiconductors. Additionally, the rise in temperature affects the material conductivity resulting in degradation of the solar cell performances. The material’s bandgap becomes narrower when the temperature increases, i.e., negative dEg/dT. Two significant reasons explain the temperature dependence of the energy gap in semiconductors–the explicit effect of the electron–phonon interaction and the implicit lattice dilation of the thermal expansion. As the temperature increases, the electrons in the cell gain more energy and become less stable, and they do not recombine with the holes before reaching the charging zone. As a result, the recombination rate of the internal carrier increases, resulting in decreasing Voc and efficiency [53, 54].

Conclusions

In this paper, we have investigated a CIGS-based solar cell using SCAPS-1D simulation. The structure involves a new ZrSxSe2−x as a nontoxic buffer layer and p-MoSe2 as an interfacial between Mo back contact and the absorber. A p-type MoSe2 layer in CIGS devices facilitates quasi-ohmic-contact development, increasing the J-V curve steepness and improving the cell performance. Additionally, the effect of thickness, Band gap values, and carrier concentration of the buffer layer of the Mo/p-MoSe2/CIGS/n-ZrSxSe2−x/ZnO structure solar cell on the photovoltaic parameters was investigated. The optimal performance parameters obtained are η 25.56%, Jsc of 37.92 mA/cm 2. Voc of 0.77 V, and FF of 85.28% at a buffer thickness of 180 nm and Eg of 1.3 eV. The increase in the ZrSxSe2−x bandgap above 1.3 eV results in a detrimental effect on the cell’s performance due to the possibility of electrons drifting toward the back contact and recombining. over, we investigated the impact of the device operating temperature on the cell performance in the range of 275 to 475 K. The observed decline in the Voc is due to the increase in the saturation current. The study presents a promising solution for Cd-free nontoxic, and efficient CIGS solar cell devices, using the ZrSxSe2−x buffer layer for the first time in CIGS solar cells.

Acknowledgements

The authors gratefully thank Marc Burgelman and his colleagues at the University of Gent, Belgium, for providing the SCAPS-1D software package used in all simulations reported in this paper.

Funding

Open access funding provided by The Science, Technology Innovation Funding Authority (STDF) in cooperation with The Egyptian Knowledge Bank (EKB).

Author information

Authors and Affiliations

• Department of Physics, School of Sciences and Engineering, The American University in Cairo, Cairo, Egypt M. Moustafa
• College of Engineering and Technology, American University of the Middle East, Kuwait, Kuwait T. Al Zoubi S. Yasin