Routes to increase performance for antimony selenide solar cells using inorganic hole transport layers.

Simple compound antimony selenide (Sb2Se3) is a promising emergent light absorber for photovoltaic applications benefiting from its outstanding photoelectric properties. Antimony selenide thin film solar cells however, are limited by low open circuit voltage due to carrier recombination at the metallic back contact interface. In this work, solar cell capacitance simulator (SCAPS) is used to interpret the effect of hole transport layers (HTL), i.e., transition metal oxides NiO and MoO x thin films on Sb2Se3 device characteristics. This reveals the critical role of NiO and MoO x in altering the energy band alignment and increasing device performance by the introduction of a high energy barrier to electrons at the rear absorber/metal interface. Close-space sublimation (CSS) and thermal evaporation (TE) techniques are applied to deposit Sb2Se3 layers in both substrate and superstrate thin film solar cells with NiO and MoO x HTLs incorporated into the device structure. The effect of the HTLs on Sb2Se3 crystallinity and solar cell performance is comprehensively studied. In superstrate device configuration, CSS-based Sb2Se3 solar cells with NiO HTL showed average improvements in open circuit voltage, short circuit current density and power conversion efficiency of 12%, 41%, and 42%, respectively, over the standard devices. Similarly, using a NiO HTL in TE-based Sb2Se3 devices improved open circuit voltage, short circuit current density and power conversion efficiency by 39%, 68%, and 92%, respectively.

1 Introduction

Antimony selenide (Sb2Se3), as a simple and low-cost compound with a direct energy band gap (∼1.18 eV), high absorption coefficient (  cm−1) and high carrier mobility (∼10 cm2/Vs, is a promising emergent light absorber for photovoltaic (PV) applications (Chen et al., 2015; Chen et al., 2017; Birkett et al., 2018). As a material, Sb2Se3 is mainly composed of (Sb4Se6) as 1-D ribbon structures, where the ribbons are strongly coupled by covalent bonds running along the c-axis with weaker Van der Waals (VdW) interactions between the ribbons. Thus, stacking of the ribbons occurs due to the weaker VdW bonds (Deringer et al., 2015). Hole mobility is enhanced in the c-axis and can reach 45 cm2/Vs along the ribbons (Black et al., 1957).

A number of studies have reported that Sb2Se3 thin films with preferred crystallographic orientation along the (hk1) direction, particularly (221), resulted in devices with higher efficiencies (Leng et al., 2014; Yuan et al., 2016; Li et al., 2017). The improved performance is often attributed to increased charge transport through the (hk1)-oriented ribbons perpendicular to the substrate and benign grain boundaries in this material (Chen et al., 2017; Williams et al., 2020). Wang and co-workers demonstrated the dependence of Sb2Se3 PV device performance on the preferred crystal orientation of the absorber (Wang et al., 2017). In that work, by optimising growth conditions, Sb2Se3 solar cells with preferred (211) and (221)-orientations on CdS and ZnO achieved higher efficiencies (5.6% and 6.0%, respectively) than those with (020) and (120)-orientations (3.2% and 4.8%, respectively). For planar Sb2Se3 solar cells in substrate orientation, a record efficiency of 6.5% has been reported with the Cd0.75Zn0.25S buffer layer being used as an alternative to CdS (Figure 1A shows standard substrate device). Meanwhile, Sb2Se3 devices with this buffer layer but in a superstrate structure (Figure 1B) have achieved an efficiency of 7.6% (Wen et al., 2018). Recently, a record substrate device efficiency of 9.2% was obtained by growing (001)-oriented Sb2Se3 nanorod arrays on sputtered molybdenum layers (Li et al., 2019). A conformal interfacial TiO2 layer was used to mitigate the migration of elemental antimony (Sb) into the CdS buffer layer, as interdiffusion has been shown to create a detrimental CdSe interlayer (Phillips et al., 2019).

FIGURE 1

Standard planar (A) substrate and (B) superstrate configuration Sb2Se3 solar cells.

In this work, thin transition metal oxides, NiO and MoO , are applied as HTLs in substrate Sb2Se3 devices to improve carrier selectivity at the back electrode by controlling inter-diffusion and formation of secondary phase materials (such as MoSe2) at the interface. Additionally, NiO and MoO HTLs are deposited on superstrate Sb2Se3 films before making Au back contacts to alter the energy band alignments at the back contact effectively producing an electron reflector, and minimising carrier recombination.

In the first part of this study, Sb2Se3 substrate/superstrate device simulations using solar cell capacitance simulator (SCAPS) are conducted in order to interpret the effect of HTLs on Sb2Se3 device characteristics (Burgelman et al., 2000). We then characterise the material properties of MoO and NiO thin films deposited at room temperature by electron beam evaporation. At this temperature it was found that NiO formed a crystalline film, unlike MoO which was amorphous. Sb2Se3 absorber films were then fabricated by close-space sublimation (CSS) and thermal evaporation (TE) techniques and incorporated into superstrate and substrate solar cell configurations. HTLs were inserted at the metal electrode/Sb2Se3 absorber interface and their effect on Sb2Se3 crystallinity and solar cell performance is comprehensively studied.

2 Experimental section

2.1 Device fabrication

The basic structure of substrate Sb2Se3 solar cells was as follows: Soda lime glass (SLG)/Mo/Sb2Se3/CdS/ZnO/ITO/Ni-Al. Mo coated soda lime glass (SLG) substrates measuring 7.5 cm2 × 2.5 cm2 were used in this study. NiO or MoO HTLs were deposited between the Mo electrode and Sb2Se3. Thin HTL films of 15 nm thickness were deposited using e-beam evaporation. 500 nm thick Sb2Se3 layers were prepared by TE of crystalline/powder Sb2Se3 source material (Alfa Aesar, 99.99%) at a deposition rate of ∼15 Å/s. The substrates were maintained at a temperature of 300°C throughout the deposition. The Sb2Se3 films were subsequently subjected to a heat treatment at 300°C for 30 min in Ar atmosphere in a tube furnace to promote recrystallisation. For the CSS Sb2Se3 films, a compact seed layer was grown at 0.05 mbar N2 for 5 min with a source temperature of 350°C, followed by a 30 min growth step at 13 mbar and a source temperature of 450°C to produce a compact and highly orientated grain structure. The substrate was then rapidly cooled with N2. An n-type CdS buffer layer (∼60 nm) was deposited by chemical bath deposition followed by DC-pulsed sputtering deposition of an i-ZnO (∼35 nm) layer plus a transparent conductive window layer ITO (∼200 nm). Front contact grids comprising Ni (∼50 nm) and Al (∼1,000 nm) were deposited through a shadow mask by e-beam evaporation. Finally, 0.16 cm2 cells were defined by mechanical scribing on each substrate.

Superstrate Sb2Se3 solar cells have the following configuration: SLG/ITO/CdS/Sb2Se3/Au with NiO or MoO HTLs deposited between the metal contact and Sb2Se3 absorber. The ITO layer was deposited by DC-pulsed sputtering and Sb2Se3 layers were grown by TE and CSS as detailed above. Finally, Au back contacts with an area of 0.07 cm2 were deposited through a shadow mask by e-beam evaporation.

2.2 Material and device characterisation

The crystal structures of Sb2Se3 were characterised by X-ray diffraction (XRD) with Cu Kα1 (1.54056 Å) radiation (Rigaku SmartLab SE). The surface morphology and cross-sectional images of Sb2Se3 films were taken by scanning electron microscopy (SEM, Tescan Mira 3 FEG-SEM). Optical spectroscopy measurements were performed using a Shimadzu UV-2600 spectrophotometer fitted with an integrating sphere. Kelvin probe force microscopy (KFPM) measurements were done using a KP Technology KP020 single point kelvin probe system fitted with a standard 2 mm Au tip.

Current-density vs. voltage (J-V) measurements of Sb2Se3 thin film solar cells were performed using an Abet Technologies solar simulator at 1-sun (100 mW/cm2) illumination equivalent to air mass 1.5 global spectrum with light power density calibrated using a Si reference cell.

2.3 Device simulation

Device simulation was carried out for both substrate and superstrate configuration Sb2Se3 solar cell using Solar Cell Capacitance Simulator (SCAPS 1-D), which is based on the solutions to Poisson’s equation and continuity equation for electrons and holes in the vertical heterostructure of multilayer thin film PV device (Burgelman et al., 2000). The input parameters of the solar cells were defined with the Sb2Se3, HTL and electron transport layer (ETL) semiconducting properties, including experimentally determined bandgaps, electron affinity, density of states (Zeng et al., 2016), mobility of charge carriers (Chen et al., 2017), acceptor/donor concentrations (Wang et al., 2015), and defect state density (Leijtens et al., 2016). Defects were introduced at the Sb2Se3/CdS interface to simulate realistic device performance.

3 Results and discussion

3.1 Simulated Sb2Se3 devices

Simulation analysis using SCAPS software was implemented to evaluate the performance of reference substrate and superstrate Sb2Se3 solar cells and those incorporating MoO and NiO as HTLs, subsequently referred to as samples Ref, MoO and NiO, respectively (see Table 1 for film properties). Figure 2 shows the J-V curves and corresponding box plots of J-V parameters of both Sb2Se3 device configurations with incorporated HTLs. Regarding the substrate devices, all device parameters are improved, with the exception of V which shows a slight decrease for devices with a HTL (down from 0.423 V for the reference device to 0.408 and 0.411 V for MoO and NiO devices, respectively). However, devices with MoO HTL show evidence of roll-over behaviour. The roll-over phenomenon, which occurs near the V in a light J-V curve, is due to Schottky energy barrier formed at the absorber/metal interface at a solar cell back contact (Eisenbarth et al., 2011; Hädrich et al., 2011). It acts as a reverse biased diode when the main junction is forward biased, blocking carrier transport for increasing forward bias, resulting in roll-over behaviour in light J-V characteristics. The baseline J in the reference device was 29.9 mA/cm2, rising to 31.2 and 31.3 mA/cm2 in MoO and NiO devices, respectively. Addition of HTL films to the reference device demonstrated a notable increase in FF for substrate devices. The FF in the reference device was 47.0%, rising to a maximum of 55.6% and 56.0% in the MoO and NiO devices, respectively. The increase in J and FF of devices with integrated HTL materials directly translates into improvements in power conversion efficiency, PCE [η = 5.9% (Ref), 6.7% (MoO ) and 7.2% (NiO)]. The current-blocking energy barrier at the back contact of the MoO substrate device could explain the lower PCE in comparison to the device with a NiO HTL. It is important to note that the results shown are not representative of the maximum conversion efficiencies that may be achieved with Sb2Se3, as we are focusing solely on the effect of the HTL, while using currently available materials parameters.

FIGURE 2

(A) J-V curves and (B) J-V parameters of simulated Sb2Se3 solar cells with different HTL materials. Roll-over behaviour is observed in the J-V curve of substrate devices with MoO HTL.

For superstrate Sb2Se3 solar cells, devices with an incorporated HTL showed an increase in J of around 8% from 31.5 mA/cm2 observed in the reference device to 33.5 and 34.0 mA/cm2 in the devices with a MoO and NiO HTL, respectively. As a result of the improvement in J , the PCE of solar cells with a HTL increased to 7.5% (MoO ) and 8.0% (NiO) from the reference value of 7.3%. Interestingly, no roll-over was seen in the J-V curve for the MoO device which could be related to the use of Au as metallic back contact rather than Mo in the substrate devices. The work function (WF) of a metal employed as a rear contact on a PV device plays an important role in facilitating hole extraction at the contact (Fleck et al., 2020). Typically, Au is reported to have a WF of 5.10 eV (Michaelson, 1977) and Mo has WFs ranging from 4.50–4.95 eV, depending on the preferred crystal orientation of the metal (Green, 1969; Michaelson, 1977; Hölzl and Schulte, 1979). To illustrate the effect of back contact metal WF on substrate/superstrate Sb2Se3 device performance, Figure 3 shows the dependence of J-V parameters on the WF of Mo and Au metals. It is apparent that the J-V parameters of all substrate devices are sensitive to variations in the value of Mo WF. In the Ref and MoO substrate devices, V decreases monotonically with Mo WF where a significant drop is observed from 0.432 V to 0.422 V at WF 4.95 eV to 0.036 V and 0.093 V at WF 4.50 eV for Ref and MoO devices, respectively. This is a clear indication of an increasing back contact barrier with decreasing Mo WF. This phenomenon has been observed experimentally in Sb2Se3 solar cells previously (Liu et al., 2014; Li et al., 2017). The V in the NiO device is less affected by the Mo WF, reducing from 0.422 V at WF 4.95 eV to 0.319 V at WF 4.50 eV. A similar trend is seen in J , FF and η parameters for the substrate devices. However, a low Mo WF of 4.50 eV causes a notable decrease in FF of the MoO device (12.3%), compared to the Ref and NiO devices (26.0% and 32.4%).

FIGURE 3

J-V parameters of simulated Sb2Se3 substrate devices with Mo back contact (varying Mo WF between 4.50–4.95 eV) and simulated Sb2Se3 superstrate devices with Au back contact (WF at 5.1 eV).

In order to understand the improvement of the device performance with the introduction of HTLs, it is necessary to consider the energy band alignment at the interfaces at the back of the PV devices. Figure 4 shows the simulated energy band diagrams of substrate and superstrate Sb2Se3 devices incorporating NiO and MoO HTLs. Due to a small electron affinity (EA = 1.46 eV (NiO), 2.05 eV (MoO )) and large band gaps (E ∼3.80 eV (NiO), 3.50 eV MoO x)) in both HTL materials, a large potential energy barrier is formed at the back contact, reflecting electrons. This barrier minimises carrier recombination at the back interfaces with Sb2Se3 and improves conductivity at the back electrode. However, it is apparent that a non-negligible hole barrier of 0.26 and 0.29 eV is formed at the MoO /Sb2Se3 interface of the substrate and superstrate devices, respectively, which can manifest as J-V roll-over behaviour seen in the simulated MoO substrate device. Thus the SCAPS simulations indicate the incorporation of a MoO or NiO HTL into substrate and superstrate configuration Sb2Se3 solar cells increases device performance compared to a standard solar cell by the introduction of a high energy barrier to electrons at the rear absorber/metal interface.

FIGURE 4

Energy level alignment for the devices in substrate (A) and superstrate (B) orientati1ons. Devices without a hole transport layer (top), with a MoO layer (middle) and a NiO layer (bottom) are shown.

3.2 Fabricated Sb2Se3 devices

100 nm thick films of MoO and NiO were deposited on SLG at room temperature to facilitate characterisation of the HTLs. Figure 5 shows surface morphology SEM images of the respective HTLs. The MoO film exhibits an amorphous, flake-like structure in comparison to a compact crystalline morphology observed in the NiO film. XRD patterns in Figure 6 confirm the amorphous and crystalline nature of the MoO and NiO films, respectively. All the diffraction peaks in the NiO thin film were identified and indexed to cubic NiO (JCPDS number 04-0835) and no diffraction peaks of other impurity phases were observed.

FIGURE 5

Top-down SEM image of a 100 nm (A) MoO film and (B) NiO films on glass. Inset: Higher magnification image of the NiO film, showing the nanostructure.

FIGURE 6

XRD pattern of 100 nm films of (A) MoOx and (B) NiO on soda lime glass (SLG). Reference XRD data for MoO2, MoO3 and NiO are shown underneath the XRD with JPDCS card ID 65-5787, 35-0609 and 04-0,835 respectively.

Supplementary Materials S1A shows the spectral transmittance and reflectance of the NiO and MoO films on SLG. Both HTLs are highly transparent in the visible and near-infrared wavelength region and their transmittance falls sharply at ultraviolet wavelengths. However, the amorphous MoO film has slightly lower transmittance/higher reflectance in the sub-600 nm wavelength region compared to the crystalline NiO film. The bandgap energy (E ) of the HTL films was calculated by extrapolation of the linear region of the Tauc plot to the x-axis, according to the relation (Tauc et al., 1966): where α is the absorption coefficient of the semiconductor material, h is Planck’s constant, ν is the frequency of the electromagnetic radiation and A is a constant of proportionality. The estimated E values of NiO and MoO films are 3.95 and 3.85 eV, respectively (see Supplementary Materials S1B). A HTL film thickness of 15 nm was incorporated into the superstrate/substrate device to ensure a conformal coating of the HTL. A HTL requires a thickness sufficient to preserve the desired material properties and not impede charge transport considerably which would detrimentally increase series resistance in the finished devices.

3.3 Superstrate devices

TE and CSS deposition techniques were employed for Sb2Se3 film growth on SLG/ITO/CdS superstrates. For TE, the SLG/ITO/CdS superstrates were heated to 300°C prior to Sb2Se3 deposition in order to promote the growth of preferred (hk1) crystal orientations while minimising (hk0) orientations (Zhou et al., 2015) (hk0) planes, specifically (120), have been found to be detrimental to carrier transport (Guo et al., 2018; Wen et al., 2018; Li et al., 2019). The (hk0)-oriented Sb2Se3 nanoribbons are stacked parallel to the ITO/SLG superstrate where conductivity is inhibited by electrically insulating VdW bonds between the stacked nanoribbons. A seed layer is used in Sb2Se3 films deposited via CSS. This seed layer has a high density of nucleation points for the second stage of growth during the CSS process, which improves uniformity, raising the average efficiency of devices (Hutter et al., 2018a). Transmittance and reflectance data for a representative TE Sb2Se3 film was used to determine the E from a Tauc plot, which gave a E value of 1.17 eV in good agreement with (Birkett et al., 2018), see Supplementary Materials S2A,B. XRD patterns for Sb2Se3 films deposited by TE and CSS are shown in Figure 7A. The peaks in both XRD patterns are sharp and well resolved indicating the polycrystalline nature of the Sb2Se3 thin films. The lattice planes are cross-referenced to JCPDS card no. 15-0861 confirming the formation of orthorhombic Sb2Se3 with space group Pbnm. Both XRD patterns show similar characteristics, exhibiting strong (211) and (221) peaks with minimal contributions from (hk0) planes. Figures 7B–E shows the top and cross-sectional SEM images of Sb2Se3 thin films deposited by TE and CSS. The different growth techniques result in contrasting Sb2Se3 film morphologies. TE produces Sb2Se3 films of uniform thickness of ∼500 nm and densely packed grains, confirming the good crystallinity of the films, consistent with the XRD results (Figures 7B,D). However, this deposition method did not form a conformal coating of the Sb2Se3 film across the entire superstrate with the presence of pinholes observed, see Supplementary Materials S3A.

FIGURE 7

(A) XRD patterns of Sb2Se3 layers deposited by TE and CSS on ITO/CdS superstrates with standard diffraction pattern for Sb2Se3 (JCPDS15-0861) included for reference and SEM images of corresponding TE (B,D) and CSS (C,E) Sb2Se3 samples.

Conversely, CSS-grown Sb2Se3 films have a rough surface morphology with exceptionally large grains in comparison to the TE films and the grains extend the full depth of the layer. Larger grains are a prerequisite for better device performance as charge mobility is faster along the Sb2Se3 ribbons than hopping between the ribbons (see Figures 7C,E). The CSS films also showed a degree of porosity but not to the extent observed in the TE films, Supplementary Materials S3B. The presence of pinholes in the Sb2Se3 films is detrimental to device performance as shunting pathways may be formed upon subsequent deposition of the Au back contact (Hutter et al., 2018b).

J-V measurements under 1-sun illumination (100 mW/cm2) were performed on Sb2Se3 devices in the standard superstrate configuration and devices incorporating MoO and NiO HTLs. The light J-V curves were fitted using a single diode model to extract the values of series (R ) and shunt (R ) resistances. Figure 8 compares the statistical distribution of the key PV parameters for these devices, where a minimum of 10 cells of each device type were measured. On average, there was a slight increase in V when a NiO HTL was incorporated into the CSS device structure. Using a NiO HTL layer increased V to 0.226 V from values of 0.201 and 0.186 V for Ref and MoO devices, respectively. The mean J of NiO cells was also enhanced to 15.94 mA/cm2 compared to Ref (11.34 mA/cm2) and MoO (10.54 mA/cm2) cells despite a slightly lower average FF in the NiO devices. This translates into a higher mean NiO CSS device efficiency of 1.01% with Ref and MoO devices achieving efficiencies of 0.71 and 0.59% respectively. Notwithstanding the higher average R (2.6 Ωcm2) and lower R (74 Ωcm2) values for NiO CSS solar cells compared to Ref (R = 2.9 Ωcm2, R = 119 Ωcm2) and MoO (R = 1.3 Ωcm2, R = 167 Ωcm2) cells, using NiO as a HTL increases performance by boosting J in CSS Sb2Se3 superstrate devices compared to the standard and MoO based devices.

FIGURE 8

J-V parameters of superstrate TE and CSS Sb2Se3 devices with incorporated MoO and NiO HTLs. □ is the average value and × is the minimum and maximum position. The three horizontal lines of each box stand for the 25%, 50%, and 75% of the reading distribution. The whisker range is determined by the standard deviation of the sampled devices. IQR is the inter-quartile range.

The average J-V parameters of TE Sb2Se3 superstrate devices followed a similar trend to those observed in the CSS devices [V : 0.209 V (Ref) → 0.214 V (MoO ) → 0.288 V (NiO), J : 1.94 mA/cm2 (MoO ) → 8.05 mA/cm2 (Ref) → 13.48 mA/cm2 (NiO) ⇒ η: 0.10% (MoO → 0.72% (Ref) → 1.38% (NiO)]. It is worth noting that the mean FF of the Ref TE cells (38.0%) was higher in relation to the cells with a HTL (27.2% MoO , 34.6% NiO). This correlates to an increase in R of 382 Ωcm2 in Ref samples from R values of 349 Ωcm2 and 154 Ωcm2 measured in MoO and NiO cells, respectively. In TE superstrate device configuration, the thin MoO film appears to form a more resistive layer compared to Ref and NiO devices (R : 10.8 Ωcm2 MoO , 4.1 Ωcm2 Ref and 6.6 Ωcm2 NiO). Thus, overall device performance in MoO based solar cells is negatively impacted by low J and high R which could be related to the amorphous nature of the MoO thin film and the presence of a current-blocking barrier at the back contact highlighted in device simulations. Despite lower FF in NiO based solar cells, device efficiencies exceed those of Ref and MoO TE devices due to improvements in V and J showing the benefit of using NiO as a HTL in superstrate Sb2Se3 solar cells.

3.4 Substrate devices

Figure 9 shows the XRD patterns of substrate Sb2Se3 thin films deposited via TE and CSS. All diffraction peaks are in good agreement with the orthorhombic Sb2Se3 (JCPDS 15-0861), which presents in the form of (hk0), (hk1) or (hk2). No diffraction peaks of other impurity phases were observed. TE Sb2Se3 films on Mo and Mo/MoO substrates show (020) and (120) peaks compared to all other Sb2Se3 films. The presence of (020) and (120) crystal orientations in thin Sb2Se3 films adversely affects PV device performance (Leng et al., 2014; Yuan et al., 2016; Li et al., 2017). However, when using a NiO HTL in TE Sb2Se3 films, it can be observed that the intensity of the diffraction peaks of Sb2Se3 is dominated by (221) and (211) crystal plane orientations. Furthermore, when using the Mo/NiO substrate, Sb2Se3 film shows an increased peak intensity for the (002) orientation. Since h and k miller indices have a zero value, it indicates that the (Sb4Se6) ribbons grow perpendicular to the substrate surface (Li et al., 2019). For CSS Sb2Se3 films, Ref and MoO samples demonstrate a higher (002) peak intensity than NiO.

FIGURE 9

XRD patterns of Sb2Se3 films deposited by TE or CSS on top of NiO, MoO and Mo-coated SLG.

Figures 10, 11 show SEM images of Sb2Se3 films on Mo-coated SLG deposited by TE and CSS methods, respectively. The top-down SEM images of the TE films (Figures 10A–C) show a difference in morphology depending on the presence of the underlying HTL. The MoO sample exhibits larger Sb2Se3 grains than the Ref sample and the presence of pinholes in both samples is patently obvious. On the other hand, the Sb2Se3 grains in the NiO sample appear more angular in nature although pinholes are still present in the film. The dissimilarity in morphology is emphasised in SEM cross-section images of the TE Sb2Se3 films (Figures 10D–F). Voids at the absorber/Mo interface are apparent in the Ref TE sample whereas the MoO sample shows a homogenous film with large grains. For the NiO sample, the Sb2Se3 grains appear column-like with no voids at the Mo interface. The top-down SEM image of all types of CSS Sb2Se3 thin films (Figures 11A–C) show significantly larger grains compared to the TE films. However, Sb2Se3 film in the Ref sample is on average thicker (∼1,000 nm) than the MoO (∼550 nm) and NiO (∼700 nm), see Figures 11D–F. The NiO sample also has a smoother surface topography.

FIGURE 10

Top-down and cross-sectional SEM images of reference substrate (A,D), MoO (B,E) and NiO (C,F) of Sb2Se3 films deposited by thermal evaporation.

FIGURE 11

Top-down and cross-sectional SEM images of reference substrate (A,D), MoO (B,E) and NiO (C,F) of Sb2Se3 films deposited by close-space sublimation.

To quantify the difference in orientations between the substrate Sb2Se3 thin films, the texture coefficient (TC) of diffraction peaks of the samples was calculated based on the following equation (Zoppi et al., 2006): where I ( is the measured peak intensity of (hkl) plane and I 0 (hkl) the intensity in the standard XRD pattern. N is the total number of reflections considered for the calculation. A diffraction peak with a relatively large TC value ( 1) indicates a preferred orientation of the grain along this direction. Figure 12 shows the TC for Sb2Se3 thin films with HTLs deposited by (A) TE and (B) CSS. It is apparent from Figure 12 that NiO HTL plays a critical role in eliminating the detrimental (hk0) planes in the TE samples and at the same time, significantly increases absorber growth in planes, i.e., (211), (221) that are perpendicular to the substrate surface. This further supports the enhanced device performance in solar cells when NiO is used as the HTL. In CSS samples, this templating effect of HTLs is not observed as no (hk0) planes are grown in the Ref and MoO samples. MoO increases the growth of favoured crystal planes including (211), (221), and (002) compared to the Ref substrate sample whereas NiO appears to inhibit the growth of the preferential planes, which may be attributed to rendering the seed layer ineffective but further study will be required to fully understand the reason.

FIGURE 12

Texture coefficient analysis from XRD patterns of Sb2Se3 films deposited via TE (A) and CSS (B) with different hole transport layers in substrate configuration. A diffraction peak with a relatively large TC value ( 1) indicates a preferred orientation of the grain along this direction.

Figure 13 shows the variation in J-V parameters measured for a minimum batch size of 10 Sb2Se3 solar cells in substrate configuration deposited by TE and CSS incorporating HTLs. The use of MoO /NiO HTLs adversely affects all device parameters in CSS-based solar cells. This can be explained by lower average R values of 55 Ωcm2 and 47 Ωcm2 determined for MoO and NiO device types, respectively, compared to 172 Ωcm2 in the Ref devices. The reason for the reduction in R of the substrate devices with a HTL is not obvious. Only working TE devices were achieved by incorporating a NiO HTL, which can be attributed to the templating effect of the NiO film which eliminated the deleterious (hk0) crystal planes and promoted the growth of preferred (211) and (221) planes. As highlighted in device simulations, the performance of substrate Sb2Se3 solar cells can be dependent on the WF of Mo back contact (see Figure 3). Mo metal typically has a WF in the range of 4.5–4.95 eV. KPFM measurements on Mo coated SLG prior to Sb2Se3 deposition determined the Mo WF to be 4.6 eV. According to simulations, device performance of Ref and MoO substrate devices is severely impacted at the observed Mo WF. Simulated NiO device performance is affected to a lesser degree.

FIGURE 13

J-V parameters of substrate TE and CSS Sb2Se3 devices with incorporated MoO and NiO HTLs. □ is the average value and × is the minimum and maximum position. The three horizontal lines of each box stand for the 25%, 50% and 75% of the reading distribution. The whisker range is determined by the standard deviation of the sampled devices. IQR is the inter-quartile range.

R values for both TE and CSS substrate Sb2Se3 devices were significantly higher than their superstrate counterparts and had a detrimental effect on overall substrate device performance. This could be related to a non-optimal sputtered ITO layer in the substrate devices with a typical sheet resistance of ∼ 35 Ω/□ (Qu et al., 2016) compared to commercially available ITO-coated glass slides used in superstrate devices with sheet resistances of 8–12 Ω/□ (Sigma Aldrich).

3.5 Simulated and fabricated device comparison

Experimentally determined device parameters, such as R and R and apparent doping density (N ) of the Sb2Se3 absorber, were incorporated into SCAPS simulations of TE/CSS Sb2Se3 devices in superstrate/substrate configurations in order to replicate the observed behaviour of the fabricated devices. For an accurate representation of the fabricated cells, the N value for the Sb2Se3 absorber in the CSS devices was set to a value previously determined for the same CSS deposition process used in this study with a Sb2Se3 absorber thickness of 1 μm (Phillips et al., 2019). An experimentally determined N value for a typical 500 nm thick TE Sb2Se3 absorber was used in TE device simulations (see Table 1 for TE/CSS Sb2Se3 film properties). Figure 14 shows device performance of the simulated TE/CSS Sb2Se3 devices with experimentally determined R , R and N values. Similar trends are observed for all device parameters of the simulated and fabricated solar cells in both device configurations indicating the simulated devices are a reasonable representation of actual Sb2Se3 solar cells (see Figures 8, 13). However, in superstrate configuration, simulations overestimate all J-V parameters, indicating factors other than R , R and N are influencing device performance. Material properties such as carrier lifetimes, defects and band tails states have been cited as having a detrimental effect on overall device performance (Chen and Tang, 2020). In that work, a number of bulk defects in Sb2Se3 were identified with energy levels within the Sb2Se3 bandgap ranging from 0.18–0.94 eV above the valence band maximum. For simulation purposes, a mid-gap donor defect (0.62 eV) was introduced for the Sb2Se3 bulk to reproduce realistic device performance (Wen et al., 2018; Ma et al., 2020). Chen and Tang (2020) also highlighted significant recombination occurring at the n-p interface which severely impacts both V and J . The presence of additional Sb2Se3 bulk defects and increased absorber/buffer interface defect concentration could account for the differences observed between the simulated and fabricated devices studied here.

TABLE 1

Device simulation parameters, d: layer thickness, E : bandgap, χ: electron affinity, ɛ/ɛ 0: dielectric constant, N : effective density of states C: conduction band (CB) V: valence band (VB), μ : carrier mobility, N : apparent doping density D: donor A: acceptor, σ : capture cross section, N : interface defect concentration, E : defect energy level relative to CB/VB and N : bulk defect concentration. Subscripts e and h are electron and hole, respectively.

Properties	MoO x	NiO	Sb2Se3	CdS	i-ZnO	ITO
d (nm)	15	15	500 (TE). 1,000 (CSS)	70	35	200
E g (eV)	3.85 a	3.95 a	1.17 a	2.72 a	3.37 b	3.72 c
χ (eV)	2.20 e	1.46 f	4.15 g	4.70 c	4.70 c	4.50 days
ɛ/ɛ 0	10.0 e	11.9 f	14.4 g	9.0 b	9.0 b	9.4 days
N C (cm−3)	2.2 × 1018e	2.2 × 1018f	2.2 × 1018g	2.1 × 1018b	1.8 × 1019b	4.0 × 1019c
N V (cm−3)	1.8 × 1019e	1.8 × 1019f	1.8 × 1019g	1.7 × 1019b	2.4 × 1018b	1.0 × 1018c
μ e (cm2/Vs.)	30 e	2.8f	100 g	160 b	200 b	30 b
μ h (cm2/Vs.)	2.5 e	2.8 f	25 g	15 b	93 b	5 b
N A/D (cm−3)	D:3 × 1016e	A:3 × 1018f	A:1 × 1014h (TE). A:1 × 1016i (CSS)	D:1 × 1017b	D:1 × 1018b	D:1 × 1021b
Defects at Sb2Se3/CdS interface (Gaussian distribution throughout interface)
N int (cm−3)			D: varied	A: varied
σ e (cm2)			10–13	10–15
σ h (cm2)			10–15	10–13
Bulk Sb2Se3 defects (Gaussian distribution throughout bulk)
N bulk (cm−3)			D: 2.6 × 1016j	A: 5.0 × 1015b
E t (eV)			0.62 j	1.20 b
σ e (cm2)			10–13	10–17
σ h (cm2)			10–15	10–13

Experimentally determined from UV-VIS, measurements.

Reference (Kanevce et al., 2015).

Reference (Erkan et al., 2016).

Reference (Kartopu et al., 2019).

Reference (Ni et al., 2019).

Reference (Casas et al., 2017).

Reference (Maurya and Singh, 2021).

Experimentally determined from capacitance-voltage C-V measurements.

Reference (Phillips et al., 2019).

Reference (Chen and Tang, 2020).

FIGURE 14

Comparison of J-V parameters of simulated TE and CSS Sb2Se3 solar cells with different HTL materials in substrate and superstrate device configurations.

In addition, it is worth noting actual superstrate devices which incorporate a MoO HTL under-perform in relation to standard simulated superstrate devices (see Figure 8). This decrease in performance is not observed in the fabricated substrate Sb2Se3 solar cells with a MoO HTL. This discrepancy can be accounted for by different processing conditions applied during deposition of substrate and superstrate devices. During deposition of Sb2Se3 layer on SLG/Mo/HTL substrate, the substrate temperature is maintained at 300°C which is sufficient to crystallise the MoO film, see Supplementary Materials S4. The crystallised MoO film consists of a mixture of MoO2, MoO3 and intermediate reduced oxide phases. The phase composition affects the electronic and optical properties of the MoO film, with MoO2 content lowering the resistivity, transmittance and bandgap (Inzani et al., 2017). Simulations also show a roll-over in the J-V curves for superstrate Sb2Se3 devices in both configurations (see Figure 15), indicating the presence of a barrier to carrier transport at the back contact seen in simulated energy band alignments as previously discussed (Figure 4).

FIGURE 15

J-V curves for simulated TE and CSS Sb2Se3 devices with HTLs in superstrate configuration. Roll-over behaviour is evident in both TE and CSS devices with MoO HTL indicating a carrier transport barrier at the back contact.

4 Conclusion

Numerical simulations of standard planar superstrate and substrate Sb2Se3 solar cells along with the effect of incorporating MoO and NiO HTLs, demonstrated an increase in device efficiency for cells with a HTL which was achieved by an increase in J for both substrate and superstrate device configurations. Both HTLs have high bandgaps and low electron affinities compared to Sb2Se3 absorber which manifests as a large barrier for electrons at the metallic back electrode and facilitates hole extraction. However, a roll-over effect was seen in the simulated J-V curve of the substrate device with MoO HTL, suggesting a current-blocking barrier at the back contact caused by non-optimal energy band alignment. Material characterisation of the HTL materials deposited by E-beam evaporation at room temperature revealed MoO formed an amorphous layer while NiO crystallised in cubic crystal orientation. 15 nm thick HTLs were incorporated into superstrate/substrate solar cells with Sb2Se3 absorbers deposited by thermal evaporation and close-space sublimation. For CSS superstrate solar cells with NiO HTL, device efficiency was enhanced by a 40% increase in J compared to reference and MoO based devices. TE superstrate cells incorporating NiO as HTL also demonstrated improved efficiencies achieved by higher V and J . In the superstrate TE cells with MoO HTL, J was severely inhibited which is attributed to MoO forming a more resistive layer due to its amorphous nature. Conversely, the presence of a MoO or NiO HTL in substrate CSS-deposited Sb2Se3 solar cells reduced device performance which is linked to lower average R observed in these cells. Optimisation of HTL thickness and/or re-optimisation of the absorber deposition could potentially alleviate this issue. Simulations reveal a connection between the WF of the Mo metal back contact and substrate device performance. For an experimentally determined Mo WF of 4.6 eV, all device J-V characteristics are significantly reduced, whereas substrate devices with NiO HTL are only marginally affected. In addition, XRD analysis of TE Sb2Se3 films with NiO HTL revealed a templating effect on Sb2Se3 crystal orientation where detrimental (020)/(120) crystal planes were eliminated and preferred (211)/(221) planes increased in intensity which resulted in increased device performance of substrate Sb2Se3 solar cells. NiO shows more promise as a HTL in Sb2Se3 PV devices, and crucially can act as a templating layer when the Sb2Se3 deposition method does not already impart the desired structure, as is often the case with TE devices.