**Influence of Surface Treatment on the Conversion Efficiency of Thin-Film a-Si:H Solar Cells on a Stainless Steel Substrate**

Wen-Cheng Ke and Shuo-Jen Lee

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/51531

## **1. Introduction**

Over the past decade, hydrogenated amorphous silicon (a-Si:H) thin-film solar cells have emerged as a viable substitute for solid-state silicon solar cells. The a-Si:H thin-film solar cells gained importance primarily due to their low production cost, and eco-friendly nature [1-3]. However, these cells have the inherent disadvantage of using glass as a substrate ma‐ terial, which makes the cells unsuitable for use in a round shaped product, are heavy and have a high material cost. Replacing the glass substrate with a stainless steel (SS) substrate makes it possible to fabricate lightweight, thin, and low-cost a-Si:H thin-film solar cells us‐ ing roll-to-roll mass production [4,5]. However, the surface morphology of a SS substrate is of poorer quality than that of the glass substrate. For the past several years flexible solar cells fabricated on a stainless steel substrate are being widely used for the building of inte‐ grated photovoltaics (BIPVs). Stainless steel has many advantages, such as low cost, high ex‐ tension, ease of preparing etc. Thus it was commonly believed that the wide application of BIPVs especially rooftop applications, would be the biggest market for flexible PV technolo‐ gy [1-4], especially since these flexible products are very light, and quick and easy to install. Indeed, BIPV products did see a rapid growth in recent years, with more companies going into production. However, the main challenge of BIPVs remains. The question of how to im‐ prove the conversion efficiency has not been resolved.

The main limitation of thin-film solar cell efficiency is the long absorption length of the long wavelength photons and the thinness of the absorbing layer. The absorption length of a-Si:H with a bandgap of 1.6 eV, for the red and infrared solar photons, exceeds 1 μm and 100 μm, respectively [6,7]. In addition, for a-Si:H the hole diffusion length is 300–400 nm, which lim‐

© 2013 Ke and Lee; licensee InTech. This is an open access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2013 Ke and Lee; licensee InTech. This is a paper distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

its the solar cell absorber layer thickness to less than the hole diffusion length [8]. This makes it exceedingly difficult to harvest these photons since the absorber thickness of a p-i-n single junction solar cell is limited to only a few hundred nanometers for efficient carrier col‐ lection. Thus, increasing the light absorption is essential for the design of thin-film solar cells [9–12]. Maximum light absorption is achieved using several different techniques including back-reflector or light trapping configurations. Enhanced light-trapping in thin-film solar cells is typically achieved by a textured oxide layer on the metal back-reflector that scatters light within the absorbing layer and increases the optical path-length of the solar photons. In our latest research [13,14], we fabricated a textured SS substrate using the photolithography method to increase the light scattering. Unfortunately, the electrical properties of a-Si:H thin-film solar cells deposited on these micro-size textured SS substrates are not stable. There are many deep etching pits on the textured SS substrate surface, resulting in a nonuniform thickness of the a-Si:H thin-film deposited on the whole textured SS substrate sur‐ face. This non-uniform coverage of the a-si:H layer on the textured 304 SS usually fails to perform its function in the thin-film solar cells. Thus, the surface morphology of the SS sub‐ strate plays an important role in achieving the stable electrical properties of the a-Si:H thinfilm solar cells.

In addition, it is well known that defects in thin-film solar cells increase the series resistance (Rs) and decrease the shunt resistance (Rsh), resulting in a decreased open voltage (VOC) and short current density (Jsc) of the thin-film solar cells. Although a number of investigations were carried out to study the properties of the defects in crystalline Si solar cells [15,16], a detailed understanding of a-Si:H thin-film solar cells is still lacking. Iron (Fe), one of the most common impurities in crystalline Si solar cell materials, has a detrimental effect on the minority carrier lifetime due to the defect states introduced by Fe and its complexes with acceptors in the band gap of silicon [17,18]. It has been suggested that diffusion of detrimen‐ tal elements, such as Fe from stainless steel, into the a-Si:H layer as a result of high tempera‐ tures during the a-Si:H processing, deteriorate the cell's efficiency. In this study, a thick (exceeding 2-μm) metal Mo buffer layer is used to reduce the diffusion of Fe impurities from the 304 SS substrate. The influence of the Fe impurities on the cell's performance was inves‐ tigated carefully. Additionally, Electro-polishing (EP) and Electrical chemical mechanical polish (ECMP) processes have been used to improve the surface roughness of the stainless steels, and make them more suitable as a substrate for a-Si:H thin-film solar cells. The EP and ECMP dissolve metal ions electrochemically by applying an anodic potential on the SS substrate surface in an aqueous electrolyte [19-22]. It should be noted that the ECMP also removes a passivation layer by the mechanical abrasion of the polishing pad and the abra‐ sives in the electrolyte. After the EP or ECMP process is finished, the SS substrate surface forms a hard and dense Cr-rich passivation layer which is expected to block the diffusion of metal impurities from the SS substrate into a-Si:H layer. We carefully investigated the con‐ version efficiency of a-Si:H thin-film solar cells based on the surface morphology and impur‐ ity diffusion of 304 SS substrate treated by the EP and ECMP processes.

## **2. Surface-treated process of 304 stainless steel substrate**

## **2.1. Electro-polishing process**

its the solar cell absorber layer thickness to less than the hole diffusion length [8]. This makes it exceedingly difficult to harvest these photons since the absorber thickness of a p-i-n single junction solar cell is limited to only a few hundred nanometers for efficient carrier col‐ lection. Thus, increasing the light absorption is essential for the design of thin-film solar cells [9–12]. Maximum light absorption is achieved using several different techniques including back-reflector or light trapping configurations. Enhanced light-trapping in thin-film solar cells is typically achieved by a textured oxide layer on the metal back-reflector that scatters light within the absorbing layer and increases the optical path-length of the solar photons. In our latest research [13,14], we fabricated a textured SS substrate using the photolithography method to increase the light scattering. Unfortunately, the electrical properties of a-Si:H thin-film solar cells deposited on these micro-size textured SS substrates are not stable. There are many deep etching pits on the textured SS substrate surface, resulting in a nonuniform thickness of the a-Si:H thin-film deposited on the whole textured SS substrate sur‐ face. This non-uniform coverage of the a-si:H layer on the textured 304 SS usually fails to perform its function in the thin-film solar cells. Thus, the surface morphology of the SS sub‐ strate plays an important role in achieving the stable electrical properties of the a-Si:H thin-

In addition, it is well known that defects in thin-film solar cells increase the series resistance (Rs) and decrease the shunt resistance (Rsh), resulting in a decreased open voltage (VOC) and short current density (Jsc) of the thin-film solar cells. Although a number of investigations were carried out to study the properties of the defects in crystalline Si solar cells [15,16], a detailed understanding of a-Si:H thin-film solar cells is still lacking. Iron (Fe), one of the most common impurities in crystalline Si solar cell materials, has a detrimental effect on the minority carrier lifetime due to the defect states introduced by Fe and its complexes with acceptors in the band gap of silicon [17,18]. It has been suggested that diffusion of detrimen‐ tal elements, such as Fe from stainless steel, into the a-Si:H layer as a result of high tempera‐ tures during the a-Si:H processing, deteriorate the cell's efficiency. In this study, a thick (exceeding 2-μm) metal Mo buffer layer is used to reduce the diffusion of Fe impurities from the 304 SS substrate. The influence of the Fe impurities on the cell's performance was inves‐ tigated carefully. Additionally, Electro-polishing (EP) and Electrical chemical mechanical polish (ECMP) processes have been used to improve the surface roughness of the stainless steels, and make them more suitable as a substrate for a-Si:H thin-film solar cells. The EP and ECMP dissolve metal ions electrochemically by applying an anodic potential on the SS substrate surface in an aqueous electrolyte [19-22]. It should be noted that the ECMP also removes a passivation layer by the mechanical abrasion of the polishing pad and the abra‐ sives in the electrolyte. After the EP or ECMP process is finished, the SS substrate surface forms a hard and dense Cr-rich passivation layer which is expected to block the diffusion of metal impurities from the SS substrate into a-Si:H layer. We carefully investigated the con‐ version efficiency of a-Si:H thin-film solar cells based on the surface morphology and impur‐

ity diffusion of 304 SS substrate treated by the EP and ECMP processes.

film solar cells.

60 Solar Cells - Research and Application Perspectives

In the 1980's, most of the research reports indicated that the EP process of the stainless steel substrate began by forming a viscous layer on the substrate surface resulting in the forma‐ tion of a passivation layer. Generally, the high spike has a higher dissolve velocity in the vis‐ cous layer providing the initial polish of the stainless steel substrate. Once the EP process has ran for a long enough time, a dense passivation layer or metal-rich oxidation layer will form at the bottom of the viscous layer on the surface of the substrate resulting in a smooth surface. The viscous layer formation mechanism in the EP pricess is a key parameter for ach‐ ieving a smooth surface stainless steel substrate. There are some conditions, such as solution convection, ion diffusion, electrical current distribution and electron migration will deter‐ mine the EP results. Thus, the applying voltage, current, gap between the anode and cath‐ ode, reaction time, cathode design must be controlled carefully. The schematic diagrams of the EP system are shown in Fig. 1. In this study, the 304 stainless steel substrate with thick‐ ness of 1 mm was clipped by the anode clamp. The stainless steel net was used as the cath‐ ode plate. The important parameter lists of the EP process are listed in Table 1. After the EP process, the mirror-like surface of the 304 stainless steel substrate can be achieved (Fig. 2).

**Figure 1.** The schematic diagram of the EP system.


**Table 1.** The key parameters of the EP process.

**Figure 2.** The image of the mirror-like 304 stainless steel substrate after the EP process.

### **2.2. Electrical chemical mechanical polish process**

The schematic diagrams of the ECMP system are shown in Fig. 3 and Fig. 4. The ECMP process consists of both electrochemical and mechanical polishing mechanisms, electrolysis and physical polishing, to repair the irregular surface morphology of the 304 SS substrate. In addition, the mechanical friction between the polishing pad and the 304 SS substrate induces an interactive force that removes the surface oxidation layer or passivation layer that forms during the electrolysis process. It should be noted that no peeling force is generated during the ECMP process, and thus there is no need to consider any mechanical deformation of the substrate. The ECMP process base on the fundamental electrochemical reaction terms as fol‐ lowing:

Anode chemical reaction:

$$\text{M}^+ + \text{OH}^- \rightarrow \text{e} \text{MOH} \tag{1}$$

Cathode chemical reaction:

$$\text{2H}^+ \rightarrow \text{\#H}\_2 \tag{2}$$

The important parameters of ECMP process are described in the list below.


The other important parameter lists of the ECMP process are listed in Table 2. In Fig. 5, the mir‐ ror-like surface of the 304 stainless steel substrate can be achieved after the ECMP process.

**Figure 3.** The schematic diagram of the ECMP system.

## **3. Effect of Mo blocking layer on the conversion efficiency of a-Si:H solar cells**

## **3.1. Experimental details**

**Figure 2.** The image of the mirror-like 304 stainless steel substrate after the EP process.

The important parameters of ECMP process are described in the list below.

The schematic diagrams of the ECMP system are shown in Fig. 3 and Fig. 4. The ECMP process consists of both electrochemical and mechanical polishing mechanisms, electrolysis and physical polishing, to repair the irregular surface morphology of the 304 SS substrate. In addition, the mechanical friction between the polishing pad and the 304 SS substrate induces an interactive force that removes the surface oxidation layer or passivation layer that forms during the electrolysis process. It should be noted that no peeling force is generated during the ECMP process, and thus there is no need to consider any mechanical deformation of the substrate. The ECMP process base on the fundamental electrochemical reaction terms as fol‐

M OH MOH *è* + - + ® (1)

<sup>2</sup> 2H H*è* <sup>+</sup> ® (2)

**2.2. Electrical chemical mechanical polish process**

62 Solar Cells - Research and Application Perspectives

lowing:

Anode chemical reaction:

Cathode chemical reaction:

In this study, a series of p-i-n solar cells were deposited on a 304 SS substrate by high fre‐ quency plasma enhanced chemical vapor deposition (HF-PECVD), in an ultrahigh-vacuum, single-chamber, load-locked system at a constant temperature of 200 °C. The

**Figure 4.** The sample holder and polishing process image of the ECMP system.


**Table 2.** The key parameters of the ECMP process.

**Figure 5.** The image of the mirror-like 304 stainless steel substrate after the ECMP process.

thickness of the p-i-n layers in the solar cell structure were 10, 300, and 20 nm respectively. In order to avoid diffusion by Fe impurities diffusing from the 304 SS substrate into Si solar cells and to study its influence on cell efficiency (see Fig. 6), we designed a series of cell structures including Ag/AZO/n-i-p/SS, Ag/AZO/n-i-p/Ag/SS, Ag/AZO/n-i-p/Ag/0.5-μmthick Mo/SS, and Ag/AZO/n-i-p/Ag/2.0-μm-thick Mo/SS. It was crucial that no aluminazinc-oxide (AZO) layer would grow between the 304 SS substrate and the a-Si:H cell so that we could clearly observe the change in cell performance as a result of preventing the diffu‐ sion of Fe impurities. All metal and transparent conductive oxide layers in the structure of the cells were deposited by radio frequency (RF) magnetron sputtering. The thickness of the back reflection metal Ag and AZO layers were 200 and 150 nm, respectively. It should be noted that the metal Mo layers, 0.5-μm-thick and 2-μm-thick were deposited on a 304 SS substrate by adjusting the sputtering power to 100 W and 500 W, respectively. The Ar gas flow rate was kept at 30 sccm, and the deposition time was kept at 30 min. In general, high sputtering power can achieve a higher metal deposition rate. The solar cell performance was determined using a calibrated AM 1.5G solar simulator under illumination, and operating at a light intensity of 100 mW/cm2 . The current-voltage (I-V) curves were obtained using the Keithley 2400 SourceMeter. The optical properties of the raw 304 SS substrate and the 304 SS substrate coated with Ag and Ag/Mo films were measured by a UV-visible(Vis)-near IR (NIR) spectrophotometer (Perkin Elmer Lambda 750s) in the 400-700 nm wavelength range. The samples were measured with a secondary ion mass spectroscope (SIMS) in order to ana‐ lyze the diffusion of Fe impurities from the 304 SS substrate into the cell structure.

**Figure 6.** Cross-sectional schematic diagrams for the samples of this study.

## **3.2. Morphological properties**

**ECMP process Parameters** Polishing time 20 min Voltage/current 7 V/1 A

40 rpm

Rotation speed (anode holder) 30 rpm

Rotati speed (Cathode head)

**Figure 5.** The image of the mirror-like 304 stainless steel substrate after the ECMP process.

thickness of the p-i-n layers in the solar cell structure were 10, 300, and 20 nm respectively. In order to avoid diffusion by Fe impurities diffusing from the 304 SS substrate into Si solar cells and to study its influence on cell efficiency (see Fig. 6), we designed a series of cell structures including Ag/AZO/n-i-p/SS, Ag/AZO/n-i-p/Ag/SS, Ag/AZO/n-i-p/Ag/0.5-μmthick Mo/SS, and Ag/AZO/n-i-p/Ag/2.0-μm-thick Mo/SS. It was crucial that no aluminazinc-oxide (AZO) layer would grow between the 304 SS substrate and the a-Si:H cell so that we could clearly observe the change in cell performance as a result of preventing the diffu‐ sion of Fe impurities. All metal and transparent conductive oxide layers in the structure of the cells were deposited by radio frequency (RF) magnetron sputtering. The thickness of the back reflection metal Ag and AZO layers were 200 and 150 nm, respectively. It should be noted that the metal Mo layers, 0.5-μm-thick and 2-μm-thick were deposited on a 304 SS substrate by adjusting the sputtering power to 100 W and 500 W, respectively. The Ar gas flow rate was kept at 30 sccm, and the deposition time was kept at 30 min. In general, high sputtering power can achieve a higher metal deposition rate. The solar cell performance was determined using a calibrated AM 1.5G solar simulator under illumination, and operating at

Keithley 2400 SourceMeter. The optical properties of the raw 304 SS substrate and the 304 SS

. The current-voltage (I-V) curves were obtained using the

**Table 2.** The key parameters of the ECMP process.

64 Solar Cells - Research and Application Perspectives

a light intensity of 100 mW/cm2

Figure 7 shows the SEM images of a metal Mo buffer layer with varying thickness on a 304 SS substrate. In Fig. 7(a), the small size of the Mo grains and the grain direction are a ran‐ dom distribution when the sputtering power is low (i.e. low deposition rate = 0.28 nm/sec). In contrast, the long ridge and the ordered arrangement of the large Mo grains was achieved by a high-power sputtering process (i.e. deposition rate = 1.11 nm/sec). In addition, the grain boundary decreased due to the increased merging of the Mo grains into the high-power sputtered sample.

## **3.3. Optical properties**

The total reflection (TR) and the diffuse reflection (DR) rates versus the wavelength curves for the raw 304 SS substrate, Ag/304 SS substrate, 0.5-μm-thick Mo/304 SS substrate and 2 μm-thick Mo/304 SS substrate are shown in Fig. 8. The TR rate is defined as the ratio of the reflection light to the incident light. When the incident light angle is zero, any reflection light with an angle larger than 8o over the incident light is called the DR rate. The TR and DR rates are very important indexes for monitoring the light-trapping in thin-film solar cells. In general, high TR and DR rates indicate that the light path can be increased in the cell structure. Thus, scattered light within the absorbing layer increases the optical path length of the solar photons. The TR and DR rates of the Mo coating on the 304 SS substrate, wheth‐ er being a 0.5-μm-thick or a 2-μm-thick Mo buffer layer were smaller than that of the raw 304 SS substrate. In addition, the TR and DR rates showed a slightly decreasing trend when the Mo buffer layer thickness increased from 0.5 μm to 2 μm. The cross-sectional SEM im‐ ages in Figs. 7(c) and (d), show the roughened surface (large grain size and deep V-shaped trench) of the 2-μm-thick Mo buffer layer, which is believed to be the main reason for the low TR and DR rates of the thick Mo buffer layer. It was also found that the TR/DR rate at a wavelength of 550 nm increased from 54.3%/7.7% of a raw 304 SS substrate to 84.2%/11.4% for the Ag coated 304 SS substrate. Previous work performed by our group showed that the surface texturing of different types of SS substrates after coating with Ag film revealed high TR and DR rates [13-14]. Coating an Ag film on the 304 SS substrate as a back reflector can help light absorption and further improve cell efficiency.

**Figure 7.** SEM images of the Mo buffer layer on a 304 SS substrate with a thickness of (a) 0.5-μm (b) 2-μm and cross sectional images of the Mo buffer layer with (c) 0.5-μm (d) 2-μm.

**Figure 8.** The TR and DR rates versus the wavelength curves of all the studied cells.

### **3.4. Current-voltage characteristics**

wavelength of 550 nm increased from 54.3%/7.7% of a raw 304 SS substrate to 84.2%/11.4% for the Ag coated 304 SS substrate. Previous work performed by our group showed that the surface texturing of different types of SS substrates after coating with Ag film revealed high TR and DR rates [13-14]. Coating an Ag film on the 304 SS substrate as a back reflector can

**Figure 7.** SEM images of the Mo buffer layer on a 304 SS substrate with a thickness of (a) 0.5-μm (b) 2-μm and cross

help light absorption and further improve cell efficiency.

66 Solar Cells - Research and Application Perspectives

sectional images of the Mo buffer layer with (c) 0.5-μm (d) 2-μm.

**Figure 8.** The TR and DR rates versus the wavelength curves of all the studied cells.

Figure 9 shows the light J-V characteristics of all studied cells. The cell structure of Ag/AZO/n-i-p/Ag/SS has poor performances with η=0.78%, FF = 29%, Voc = 0.44 V and Jsc = 6.18 mA/cm2 . This poor cell efficiency may be due to several reasons, including incomplete absorption of incident light, and the presence of intrinsic or processing-induced defects. The light absorption of thin film solar cells can be enhanced by using the surface textured 304 SS substrate as a back-reflector. Since the DR and TR rates of the raw 304 SS substrate is low, a 200 nm thick Ag film was coated on the 304 SS substrate as a back-reflector. In Fig. 8, the increased TR and DR rates of the Ag coated 304 SS substrate was helpful for increasing the Voc and Jsc of the a-Si:H thin film solar cell. However, the cell conversion efficiency of 1.3% is still too low for a-Si:H thin film solar cells, and the increased cell efficiency by using a metal Ag back-reflector is limited. On the other hand, the Fe impurities or iron-boron complex are deep-level defects in Si [23,24], and are known to have serious detrimental effects on the effi‐ ciency of crystalline Si solar cells. Thus, we suggest that these Fe impurities diffused from the 304 SS substrate into the a-Si:H thin film solar cells also form deep-level defects that de‐ teriorates the cell's performance. A detailed and comprehensive understanding is necessary to address this process. In addition, a diffusion blocking layer is necessary to prevent the Fe impurities diffused from the 304 SS substrate to enter the a-Si:H thin film solar cells.

**Figure 9.** The J-V characteristics for all the studied cells.

### **3.5. Impurity diffusion in a-Si:H films**

In order to study this Fe diffusion, we performed a SIMS analysis. Figure 10 shows the SIMS depth profile of the p-i-n a-Si:H thin film cells on the Ag/ 304 SS substrate and the Ag/2-μmthick Mo/304 SS substrate. In Fig. 10(a), the Fe atom was detected in the whole p-i-n cell structure. In particular, the maximum signal intensity of the Fe atom was found at a depth of 500 nm in the Ag/AZO/n-i-p/Ag/SS substrate sample. The driving force of the Fe diffusion from the 304 SS substrate into the p-i-n a-Si:H solar cells was believed to be due to the high temperature of the PECVD a-Si:H deposition process. The 200 nm Ag back-reflector can't prevent the Fe diffusion from the 304 SS substrate. In Fig. 10(b), the Fe atom signal intensity from the 304 SS substrate decreased with the increase in the thickness of the Mo buffer layer. In addition, almost no Fe atoms were detected in the p-i-n a-Si:H solar cells. It should be mentioned that few Fe atoms can diffuse through the 0.5-μm-thick Mo film into the a-Si:H layer. In Fig. 10(c), the cross-sectional SEM image indicates that the low sputtering rate of the Mo film has a dense columnar structure. High density, thin Mo films may be useful to suppress the diffusion of Fe atoms. However, we must consider the time-cost estimation. A 2-μm thick Mo film deposited at the low sputtering rate (0.28 nm/sec) requires 2 hrs, and at the high sputtering rate (1.11 nm/sec) it only requires 30 min. In addition, a 2-μm thick Mo film deposited at the high sputtering rate can effectively suppress the diffusion of Fe atoms into the a-Si:H layer. Thus, the optimal deposition condition for the Mo buffer layer is be‐ lieved to be a thicker layer deposited at a high sputtering rate.

**Figure 10.** The SIMS depth profile of a-Si:H solar cells grown on (a) an Ag/304 SS substrate and (b) an Ag/2-μm Mo/304 SS substrate.

#### **3.6. a-Si:H thin-film solar cells performance**

Table 3 summarizes the relative cell performance of all studied cells. The best cell structure (i.e. Ag/AZO/n-i-p/Ag/2.0-μm-thick Mo/ SS) has a performance of η=4.12%, FF = 55%, Voc = 0.78V and Jsc = 9.54 mA/cm<sup>2</sup> . The Rsh was 119.1 Ω and 290.6 Ω for a-Si:H solar cell on the Ag coated 304 SS substrate and the Ag/2-μm-thick Mo film/304 SS substrate, respectively. In addition, the ser‐ ies resistance (Rs) of the Ag/AZO/n-i-p/Ag/ SS cell decreased from 53.3 Ω to 16.7 Ω for the Ag/AZO/n-i-p/Ag/2.0-μm-thick Mo/ SS cell. There are many reasons for the series resistance, including the resistance of the semiconductor bulk, contacts, interconnections, etc. Shunt re‐ sistance is believed to be caused mainly by lattice defects or leakage current at the edge of the solar cell. Based on the SIMS results, we believe that the high conversion efficiency was due to the effective prevention of the diffusion of Fe impurities. Fe impurities in the Si thin films will form deep-level defects that degenerate the junction interface in each layer of the cell structure thereby decreasing the Rsh of the solar cell. In addition, the Fe induced deep-level defects will increase the resistivity of the intrinsic a-Si:H layer which may raise the Rs. The high Voc and Jsc were due to an increase of Rsh and a decrease of Rs as a result of coating a thick Mo buffer lay‐ er on the 304 SS substrate. This effectively prevented Fe impurities from diffusing into the a-Si:H solar cells thereby improving the conversion efficiency of the solar cells. The improvement of the conversion efficiency of a-Si:H thin film solar cells on a 304 SS substrate by using a Mo buffer layer is still an open question. Some possible reasons might explain the im‐ provement in cell efficiency by using a thick Mo buffer layer. For example, a thick Mo buffer layer decreases the surface roughness of the raw 304 SS substrate. The decrease in interface de‐ fect density of the a-Si:H thin-film solar cells due to the improved interface flatness results in an increased conversion efficiency as the Mo buffer layer thickness increases. However, we have to keep in mind that the thermal expansion coefficients of Si, Mo and stainless steel are 2.5×10-6/ K, 5.2×10-6/K and 1.4×10-5/K, respectively. Consequently, the a-Si:H layer grown directly on the stainless steel substrate will induce structural defects because of the severe mismatch in ther‐ mal expansion coefficients. The low thermal expansion coefficient of Mo film onto a stainless steel substrate can minimize the mismatch in thermal expansion coefficients between the a-Si:H layer and the 304 SS substrate. The density of the structural defects in the a-Si:H thin-film solar cells can be reduced, resulting in a high conversion efficiency of thin-film a-Si:H solar cells on the 304 SS substrate. More detailed experiments must be carried out in the future to al‐ low us to understand the precise mechanism involved in improving cell efficiency by using a Mo buffer layer.


**Table 3.** AM 1.5G output parameters of the a-Si:Hthin-film solar cells.

## **4. Surface treatment of the 304 stainless steel substrte to improve the conversion efficiency of a-Si:H solar cells**

### **4.1. Experimental details**

from the 304 SS substrate into the p-i-n a-Si:H solar cells was believed to be due to the high temperature of the PECVD a-Si:H deposition process. The 200 nm Ag back-reflector can't prevent the Fe diffusion from the 304 SS substrate. In Fig. 10(b), the Fe atom signal intensity from the 304 SS substrate decreased with the increase in the thickness of the Mo buffer layer. In addition, almost no Fe atoms were detected in the p-i-n a-Si:H solar cells. It should be mentioned that few Fe atoms can diffuse through the 0.5-μm-thick Mo film into the a-Si:H layer. In Fig. 10(c), the cross-sectional SEM image indicates that the low sputtering rate of the Mo film has a dense columnar structure. High density, thin Mo films may be useful to suppress the diffusion of Fe atoms. However, we must consider the time-cost estimation. A 2-μm thick Mo film deposited at the low sputtering rate (0.28 nm/sec) requires 2 hrs, and at the high sputtering rate (1.11 nm/sec) it only requires 30 min. In addition, a 2-μm thick Mo film deposited at the high sputtering rate can effectively suppress the diffusion of Fe atoms into the a-Si:H layer. Thus, the optimal deposition condition for the Mo buffer layer is be‐

**Figure 10.** The SIMS depth profile of a-Si:H solar cells grown on (a) an Ag/304 SS substrate and (b) an Ag/2-μm

Table 3 summarizes the relative cell performance of all studied cells. The best cell structure (i.e. Ag/AZO/n-i-p/Ag/2.0-μm-thick Mo/ SS) has a performance of η=4.12%, FF = 55%, Voc = 0.78V

SS substrate and the Ag/2-μm-thick Mo film/304 SS substrate, respectively. In addition, the ser‐ ies resistance (Rs) of the Ag/AZO/n-i-p/Ag/ SS cell decreased from 53.3 Ω to 16.7 Ω for the Ag/AZO/n-i-p/Ag/2.0-μm-thick Mo/ SS cell. There are many reasons for the series resistance, including the resistance of the semiconductor bulk, contacts, interconnections, etc. Shunt re‐ sistance is believed to be caused mainly by lattice defects or leakage current at the edge of the solar cell. Based on the SIMS results, we believe that the high conversion efficiency was due to the effective prevention of the diffusion of Fe impurities. Fe impurities in the Si thin films will form deep-level defects that degenerate the junction interface in each layer of the cell structure thereby decreasing the Rsh of the solar cell. In addition, the Fe induced deep-level defects will increase the resistivity of the intrinsic a-Si:H layer which may raise the Rs. The high Voc and Jsc

. The Rsh was 119.1 Ω and 290.6 Ω for a-Si:H solar cell on the Ag coated 304

lieved to be a thicker layer deposited at a high sputtering rate.

68 Solar Cells - Research and Application Perspectives

Mo/304 SS substrate.

and Jsc = 9.54 mA/cm<sup>2</sup>

**3.6. a-Si:H thin-film solar cells performance**

The a-Si:H thin-film solar cells were grown on 1-mm-thick 304 SS substrates by high-fre‐ quency plasma enhanced chemical vapor deposition (HF-PECVD) in an ultrahigh-vacuum, single-chamber, load-locked system at a constant temperature of 200 °C. The cell structures (see Fig. 11), consisting of Ag finger electrode/AZO/n-i-p/Ag/304 SS substrates were used to study the cells' performance on the EP and ECMP processed 304 SS substrate. The 60 nm Ag layer was deposited on the 304 SS substrate as a back reflector layer and back contact by ra‐ dio frequency (RF) sputtering. The thickness of the n-i-p layers in the solar cell structure were 20, 300 and 10 nm, respectively. The 150-nm-thick AZO films were then grown on the a-Si:H active layer by RF-sputtering. The EP process is an anodic dissolution of the surface polishing process. The 304 SS substrate was immersed in the electrolyte, and then a current of 4.5 A and voltage of 3.3 V was applied for 1 minute. For the ECMP process, the 304 SS substrate was treated with 3%wt NaNO3 solution and 1 μm diameter alumina powders. The formation of the passivation layer on the 304 SS substrate during the initial ECMP process led to spikes on the sample surface that were readily removed by powder polishing. The ECMP process required roughly 5 minutes to achieve a mirror-like surface on the 304 SS substrate. The solar cell performance was measured using a calibrated AM 1.5G solar simu‐ lator under illumination, and operating at a light intensity of 100 mW/cm2 . The current den‐ sity-voltage (J-V) curves were obtained using the Keithley 2400 Source Meter. The optical properties of the untreated 304 SS substrate and the EM/ECMP processed 304 SS substrate were measured by a UV-visible(Vis)-near IR (NIR) spectrophotometer (Perkin Elmer Lamb‐ da 750s) in the 300-1000 nm wavelength range. The samples were measured with a secon‐ dary ion mass spectroscope (SIMS) in order to analyze the diffusion of iron (Fe) and chromium (Cr) impurities from the 304 SS substrate into the cell structure.

**Figure 11.** Cross-sectional schematic diagrams for the samples of this study.

### **4.2. Morphological properties**

Fig. 12 shows the SEM and optical microscopy images of the untreated 304 SS substrate and the 304 SS substrate treated by EP and ECMP process. Fig. 12(a), shows many voids and scratches on the untreated 340 SS substrate surface. The surface roughness measured by 3D confocal microscopy, shows an average roughness (Ra) of ~ 330 nm for a surface area of 300×300 μm2 for the untreated 304 SS substrate. Fig. 12(b), shows that the voids and scratch‐ es on the surface of the 304 SS substrate can be readily removed after the EP process. The smooth surface of the 304 SS substrate by means of the EP process was accomplished using the anode electrolysis mechanism that places the 304 SS substrate in the electrolyte and then applies the optimal voltage and current between the anode and the cathode electrode. In the 1980's, most of the research reports indicated that the EP process of the stainless steel sub‐ strate began by forming a viscous layer on the substrate surface resulting in the formation of a passivation layer. Generally, the high spike has a higher dissolve velocity in the viscous layer providing the initial polish of the stainless steel substrate. Once the EP process has ran for a long enough time, a dense passivation layer or metal-rich oxidation layer will form at the bottom of the viscous layer on the surface of the substrate resulting in a smooth surface. However, the shortcoming of the EP process is that some small scratches and pin holes may remain that it cannot remove completely. Thus, this study used the ECMP process in an at‐ tempt to achieve a smoother surface of the 304 SS substrate. The ECMP process consists of both electrochemical and mechanical polishing mechanisms, electrolysis and physical pol‐ ishing, to repair the irregular surface morphology of the 304 SS substrate. In addition, the mechanical friction between the polishing pad and the 304 SS substrate induces an interac‐ tive force that removes the surface oxidation layer or passivation layer that forms during the electrolysis process. It should be noted that no peeling force is generated during the ECMP process, and thus there is no need to consider any mechanical deformation of the substrate. Our experimental results indicated that the surface roughness of the 304 SS substrate can be reduced to 22 nm after having been treated by the ECMP process, as shown in Fig. 12(c).

**Figure 12.** SEM and OM images of the (a) untreated 304 SS substrate, (b) EP, and (c) ECMP processed 304 SS sub‐ strate.

## **4.3. Optical properties**

a-Si:H active layer by RF-sputtering. The EP process is an anodic dissolution of the surface polishing process. The 304 SS substrate was immersed in the electrolyte, and then a current of 4.5 A and voltage of 3.3 V was applied for 1 minute. For the ECMP process, the 304 SS substrate was treated with 3%wt NaNO3 solution and 1 μm diameter alumina powders. The formation of the passivation layer on the 304 SS substrate during the initial ECMP process led to spikes on the sample surface that were readily removed by powder polishing. The ECMP process required roughly 5 minutes to achieve a mirror-like surface on the 304 SS substrate. The solar cell performance was measured using a calibrated AM 1.5G solar simu‐

sity-voltage (J-V) curves were obtained using the Keithley 2400 Source Meter. The optical properties of the untreated 304 SS substrate and the EM/ECMP processed 304 SS substrate were measured by a UV-visible(Vis)-near IR (NIR) spectrophotometer (Perkin Elmer Lamb‐ da 750s) in the 300-1000 nm wavelength range. The samples were measured with a secon‐ dary ion mass spectroscope (SIMS) in order to analyze the diffusion of iron (Fe) and

Fig. 12 shows the SEM and optical microscopy images of the untreated 304 SS substrate and the 304 SS substrate treated by EP and ECMP process. Fig. 12(a), shows many voids and scratches on the untreated 340 SS substrate surface. The surface roughness measured by 3D confocal microscopy, shows an average roughness (Ra) of ~ 330 nm for a surface area of 300×300 μm2 for the untreated 304 SS substrate. Fig. 12(b), shows that the voids and scratch‐ es on the surface of the 304 SS substrate can be readily removed after the EP process. The smooth surface of the 304 SS substrate by means of the EP process was accomplished using the anode electrolysis mechanism that places the 304 SS substrate in the electrolyte and then applies the optimal voltage and current between the anode and the cathode electrode. In the 1980's, most of the research reports indicated that the EP process of the stainless steel sub‐ strate began by forming a viscous layer on the substrate surface resulting in the formation of a passivation layer. Generally, the high spike has a higher dissolve velocity in the viscous layer providing the initial polish of the stainless steel substrate. Once the EP process has ran for a long enough time, a dense passivation layer or metal-rich oxidation layer will form at

. The current den‐

lator under illumination, and operating at a light intensity of 100 mW/cm2

chromium (Cr) impurities from the 304 SS substrate into the cell structure.

**Figure 11.** Cross-sectional schematic diagrams for the samples of this study.

**4.2. Morphological properties**

70 Solar Cells - Research and Application Perspectives

The effect of the surface morphology of the 304 SS substrate, treated by the EP and ECMP processes, on the light reflection properties were measured with a UV-visible-nIR spectrom‐ eter in the wavelength range of 300-1000 nm. As shown in Fig. 13, the average total reflec‐ tion (TR) rate in the wavelength range of 300-1000 nm increased from 65.6 % for the untreated 304 SS substrate to 74.2 % and 73.7% respectively for the EP and ECMP processed 304 SS substrate. The average diffuse reflection (DR) rate was 3.8% for the untreated 304 SS substrate and decreased to 0.9% and 0.8% for the EP and ECMP processed 304 SS substrate, respectively. The high TR rate and low DR rate of the EP and ECMP processed 304 SS sub‐ strate were consistent with the SEM results. Thus, the EP and ECMP processes can achieve a smooth surface of the 304 SS substrate.

### **4.4. Current-voltage characteristics**

Figure 14 shows the light current density and voltage (J-V) characteristics of all cells studied. The cell structure of the Ag finger electrode/AZO/n-i-p/Ag/untreated 304 SS substrate per‐ formed poorly with η=1.7 %, FF = 0.39, Jsc = 5.1 mA/cm<sup>2</sup> , and Voc = 0.79 V. Compared to the elec‐ trical properties of the a-Si:H thin-film solar cells grown on the glass substrate, the Jsc (i.e. ~ 15 mA/cm2 ) is a key electrical property that needs to be improve in order to achieve a high con‐ version efficiency of the a-Si:H thin-film solar cells on the 304 SS substrate. We first considered the incomplete absorption of incident light by the a-Si:H layer. It is known that the light ab‐ sorption of thin-film solar cells can be enhanced by using a high reflectance metal as a back-re‐ flector. Fig. 13 shows that the increased TR rate of the EP and ECMP processed 304 SS substrate was helpful to increase the JSC of the a-Si:H thin-film solar cells. Compared with the a-Si:H thinfilm solar cells on the untreated 304 SS substrate, the Jsc can be increased by about 35% and 43% for cells on EP and ECMP processed 304 SS substrates, respectively.

**Figure 13.** The TR and DR rates versus the wavelength curves of all the studied SS substrate samples.

It is worth noting that the difference of the average TR rates between the EP and ECMP processed 304 SS substrate was about 0.5%. However, the Jsc increased to 7.9 mA/cm2 for cells on the EP processed 304 SS substrate, and to 9.0 mA/cm2 for cells on the ECMP process‐ ed 304 SS substrate. Thus, another way to improve the Jsc of cells was associated with the intrinsic or processing-induced defects. In our previous study, we found that the Fe impuri‐ ties that diffused from the 304 SS substrate into the a-Si:H layer deteriorated the cells' per‐ formance. The Fe impurities of the iron-boron complex are deep-level defects in Si, and are known to have serious detrimental effects on the efficiency of crystalline Si solar cells [25,26]. Thus, we believe that the impurities diffused from the 304 SS substrate into the a-Si:H thin-film solar cells form deep-level defects that deteriorate the cell's performance. A detailed and comprehensive understanding of the process is required to address this issue.

### **4.5. Impurity diffusion in a-Si:H films**

In order to study the diffusion of impurities from the 304 SS substrate, we performed a SIMS depth profile analysis of the cell structure. Fig. 15 shows the SIMS depth profile of the p-i-n a-Si:H thin-film solar cells on the Ag/untreated 304 SS substrate, the Ag/EP processed 304 SS substrate and the Ag/ECMP processed 304 SS substrate. In Fig. 15(a), the Fe and Cr atoms were detected in the whole p-i-n cell structure. The driving force of the Fe and Cr atoms dif‐ fusion from 304 SS substrate into the p-i-n a-Si:H solar cells was believed to be due to the high temperature of the PECVD a-Si:H layer deposition process. Fig. 15(b) shows that the intensity of the Fe and Cr atoms signals from the 304 SS substrate can be decreased by three orders of magnitude in the p-i-n a-Si:H layers. The dense Cr-rich passivation layer formed on the surface of the 304 SS substrate by the EP process is believed to act as a blocking layer to suppress the continual diffusion of Fe and Cr atoms from the 304 SS substrate during the high-temperature PECVD deposition process. The passivation layer can be a thin layer of oxidized metal that forms during the EP process. However, there remain some small scratches on the EP processed 304 SS substrate (see Fig. 12(b)) that provide the impurities diffusion with a channel into the a-Si:H layer. Thus, a small amount of Cr and Fe impurities can still diffuse into the a-Si:H layer. On the other hand, Fig. 15(c) shows that almost no Cr atoms are detected in the p-i-n a-Si:H layer. The ECMP process includes a mechanical pol‐ ishing mechanism, and the mechanical friction between the polishing pad and the 304 SS substrate induces an interactive force that removes the Cr-rich passivation layer which formed during the electrolysis. Thus, we believe that the thickness of the Cr-rich passivation layer or oxidation layer on the ECMP processed 304 SS substrate is thinner than that of the EP processed 304 SS substrate. This thinner Cr-rich passivation layer may allow Fe atoms to diffuse into the a-Si:H layer. However, the dense and hard Cr-rich passivation layer on the 304 SS substrate which is generated during the electrolysis process will suppress the Cr atoms of the 304 SS substrate diffused into the a-Si:H layer.

**Figure 14.** The J-V characteristics for all the studies cells.

version efficiency of the a-Si:H thin-film solar cells on the 304 SS substrate. We first considered the incomplete absorption of incident light by the a-Si:H layer. It is known that the light ab‐ sorption of thin-film solar cells can be enhanced by using a high reflectance metal as a back-re‐ flector. Fig. 13 shows that the increased TR rate of the EP and ECMP processed 304 SS substrate was helpful to increase the JSC of the a-Si:H thin-film solar cells. Compared with the a-Si:H thinfilm solar cells on the untreated 304 SS substrate, the Jsc can be increased by about 35% and 43%

for cells on EP and ECMP processed 304 SS substrates, respectively.

72 Solar Cells - Research and Application Perspectives

**Figure 13.** The TR and DR rates versus the wavelength curves of all the studied SS substrate samples.

**4.5. Impurity diffusion in a-Si:H films**

It is worth noting that the difference of the average TR rates between the EP and ECMP processed 304 SS substrate was about 0.5%. However, the Jsc increased to 7.9 mA/cm2 for cells on the EP processed 304 SS substrate, and to 9.0 mA/cm2 for cells on the ECMP process‐ ed 304 SS substrate. Thus, another way to improve the Jsc of cells was associated with the intrinsic or processing-induced defects. In our previous study, we found that the Fe impuri‐ ties that diffused from the 304 SS substrate into the a-Si:H layer deteriorated the cells' per‐ formance. The Fe impurities of the iron-boron complex are deep-level defects in Si, and are known to have serious detrimental effects on the efficiency of crystalline Si solar cells [25,26]. Thus, we believe that the impurities diffused from the 304 SS substrate into the a-Si:H thin-film solar cells form deep-level defects that deteriorate the cell's performance. A detailed and comprehensive understanding of the process is required to address this issue.

In order to study the diffusion of impurities from the 304 SS substrate, we performed a SIMS depth profile analysis of the cell structure. Fig. 15 shows the SIMS depth profile of the p-i-n a-Si:H thin-film solar cells on the Ag/untreated 304 SS substrate, the Ag/EP processed 304 SS substrate and the Ag/ECMP processed 304 SS substrate. In Fig. 15(a), the Fe and Cr atoms

### **4.6. a-Si:H thin-film solar cells performance**

Table 4 summarizes the relative cell performance of all cells studied. The best cell for the a-Si:H grown on the ECMP processed 304 SS substrate had a performance of η =3.7 %, FF = 0.52, Jsc = 9.0 mA/cm2 , and Voc = 0.78 V. The Rsh was 923 Ω and 1269 Ω for a-Si:H solar cells grown on the untreated 304 SS substrate and ECMP processed 304 SS substrate, respectively. In addition, the series resistance (Rs) of the untreated 304 SS substrate cell decreased from 136 Ω to 72 Ω for the ECMP processed 304 SS substrate cell. There are many reasons for the series resistance, including the resistance of the semiconductor bulk, contacts, interconnec‐ tions, etc. The shunt resistance is believed to be caused mainly by lattice defects or leakage current at the edge of the solar cell. Based on the SIMS results, we believe that the high con‐ version efficiency was due to the effective prevention of the diffusion of Cr and Fe impuri‐ ties. The Cr and Fe impurities in the Si thin films form deep-level defects that degenerate the junction interface in each layer of the cell structure thereby decreasing the Rsh of the solar cell. In addition, the Cr and Fe induced deep-level defects increased the resistivity of the in‐ trinsic a-Si:H layer which could increase the Rs. The high Jsc was due to an increase in Rsh and a decrease in Rs as a result of EP and ECMP processed 304 SS substrate. These processes effectively prevented Cr and Fe impurities from diffusing into the a-Si:H solar cells thereby improving the conversion efficiency of the solar cells.

**Figure 15.** The SIMS depth profile of a-Si:H thin-film solar cells grown on (a) untreated 304 SS substrate, (b) EP and (c) ECMP processed 304 SS substrate.

In addition, the surface morphology and the diffusion of impurities into the 304 SS substrate influence the electrical performance of the a-Si:H thin-film solar cells. We must also keep in mind that the thermal expansion coefficients of Si, Ag and stainless steel are 2.5×10-6/K, 1.95×10-5/K and 1.4×10-5/K, respectively. Consequently, the a-Si:H layer grown directly on the

Ag/stainless steel substrate will induce structural defects as a result of the severe mismatch in thermal expansion coefficients. In order to improve the thermal expansion mismatch is‐ sue, we inserted the ZnO:Al buffer layer between the a-Si:H layer and Ag/304 SS substrate. The thermal expansion coefficient of ZnO:Al is 5.3×10-6/K. The low thermal expansion coeffi‐ cient of ZnO:Al film onto a Ag/304 SS substrate can minimize the mismatch in thermal ex‐ pansion coefficients between the a-Si:H layer and the 304 SS substrate. This reduces the density of the structural defects in the a-Si:H thin-film solar cells, resulting in a higher con‐ version efficiency of 4.5 % for a-Si:H thin-film solar cells on the ECMP processed 304 SS sub‐ strate. More detailed experiments including putting n-type a-Si:H down on the SS substrate side, adjusting the thickness of the ZnO:Al & Ag back-reflector must be carried out in the future to allow us to further improve cell efficiency.


**Table 4.** AM 1.5G output parameters of the a-Si:H thin-film solar cells.

## **5. Conclusions**

grown on the untreated 304 SS substrate and ECMP processed 304 SS substrate, respectively. In addition, the series resistance (Rs) of the untreated 304 SS substrate cell decreased from 136 Ω to 72 Ω for the ECMP processed 304 SS substrate cell. There are many reasons for the series resistance, including the resistance of the semiconductor bulk, contacts, interconnec‐ tions, etc. The shunt resistance is believed to be caused mainly by lattice defects or leakage current at the edge of the solar cell. Based on the SIMS results, we believe that the high con‐ version efficiency was due to the effective prevention of the diffusion of Cr and Fe impuri‐ ties. The Cr and Fe impurities in the Si thin films form deep-level defects that degenerate the junction interface in each layer of the cell structure thereby decreasing the Rsh of the solar cell. In addition, the Cr and Fe induced deep-level defects increased the resistivity of the in‐ trinsic a-Si:H layer which could increase the Rs. The high Jsc was due to an increase in Rsh and a decrease in Rs as a result of EP and ECMP processed 304 SS substrate. These processes effectively prevented Cr and Fe impurities from diffusing into the a-Si:H solar cells thereby

**Figure 15.** The SIMS depth profile of a-Si:H thin-film solar cells grown on (a) untreated 304 SS substrate, (b) EP and (c)

In addition, the surface morphology and the diffusion of impurities into the 304 SS substrate influence the electrical performance of the a-Si:H thin-film solar cells. We must also keep in mind that the thermal expansion coefficients of Si, Ag and stainless steel are 2.5×10-6/K, 1.95×10-5/K and 1.4×10-5/K, respectively. Consequently, the a-Si:H layer grown directly on the

improving the conversion efficiency of the solar cells.

74 Solar Cells - Research and Application Perspectives

ECMP processed 304 SS substrate.

This study presented the influence of the thickness of the metal Mo buffer layer on a 304 SS substrate on the performance of a-Si:H solar cells. The SIMS result showed that the Fe im‐ purities can be blocked effectively by increasing the thickness of the Mo buffer layer to more than 2 μm. The increased Voc and Jsc of a-Si:H solar cells on a Ag/Mo/304 SS substrate was due to an increased Rsh and a decreased Rs which related to the reduction of the Fe deeplevel defects density. The Mo buffer layer functioning as an Fe impurities blocking layer plays an important role in improving the a-Si:H thin-film solar cells on a 304 SS substrate.

EP and ECMP surface treatment techniques were also used to smoothen the 304 SS substrate surface. A decreased surface roughness of untreated 304 SS substrate as a result of being subjected to the EP or ECMP process increased the TR rate. We believe that the high TR rate due to the smooth surface of the 304 SS substrate after undergoing the EP or ECMP process increased the short current density and as a result increased the cell conversion efficiency. In addition, the SIMS analysis indicated that the diffusion of Fe and Cr impurities from the 304 SS substrate into the a-Si:H solar cell can be suppressed by using the EP process. We sug‐ gested that due to the dense and hard Cr-rich passivation layer that was formed on the ECMP processed 304 SS substrate, the Cr impurity was nearly entirely prevented from dif‐ fusing into the a-Si:H layer, resulting in a decreased RS and increase Rsh of the cell. The smooth surface and the low level of diffusion of impurities of the ECMP processed 304 SS substrate play an important role in improving the conversion efficiency of the a-Si:H thinfilm solar cells.

### **Acknowledgements**

The authors gratefully acknowledge the financial support from the National Science Council of Taiwan, R.O.C. under contract nos. NSC-98-2112-M155-001-MY3 and NSC-99-2221- E-155-065.

## **Author details**

Wen-Cheng Ke and Shuo-Jen Lee

Department of Mechanical Engineering, Yuan Ze University, Chung-Li, 320, Taiwan, R.O.C.

## **References**


[12] Müller, J., Rech, B., Springer, J., & Vanecek, M. (2004). *Sol. Energy, 77 917*.

substrate play an important role in improving the conversion efficiency of the a-Si:H thin-

The authors gratefully acknowledge the financial support from the National Science Council of Taiwan, R.O.C. under contract nos. NSC-98-2112-M155-001-MY3 and NSC-99-2221-

Department of Mechanical Engineering, Yuan Ze University, Chung-Li, 320, Taiwan, R.O.C.

[1] Hsu, . M., Tathireddy, P., Rieth, L., Normann, A. R., & Solzbacher, F. (2007). *Thin Sol‐*

[2] Fan, Q. H., Chen, C., Liao, X., Xiang, X., Zhang, S., Ingler, W., Adiga, N., Hu, Z., Cao,

[3] Huran, J., Hotovy, I., Pezoltd, J., Balaykin, N. I., & Kobzev, A. P. (2006). *Thin Solid*

[4] Deng, X. M., Liao, X. B., Han, S. J., Povolny, H., & Agarwal, P. (2000). *Sol. Energy Ma‐*

[6] Ferlanto, A. S., Ferreira, G. M., Pearce, J. M., Wronski, C. R., Collins, R. W., Deng, X.,

[9] Anna, J. A., Selvan, A. E., Delahoy, S., Guo, , & Li, Y. M. (2006). *Sol. Energy Mater. Sol.*

[11] Sőderstrőm, T., Haug, F. J., Terrazzoni-Daudrix, V., & Ballif, C. (2008). *J. Appl. Phys.*

[5] Kolodziej, A., Krewniak, P., & Nowak, S. (2003). *Opto-electronics Review, 11 281*.

X., Du, W., & Deng, X. M. (2010). *Sol. Energy Mater. Sol. Cells 94 1300*.

film solar cells.

E-155-065.

**Author details**

**References**

Wen-Cheng Ke and Shuo-Jen Lee

*id Films 516 34.*

*Films 515 651*.

*Cells, 90 3371*.

*103 114509*.

*ter. Sol. Cells 62 89*.

& Ganguly, G. (2002). *J. Appl. Phys. 92 2424*.

[7] Zhou, D., & Biswas, R. (2008). *J. Appl. Phys. 103 093102*.

[8] Curtin, B., Biswas, R., & Dalal, V. (2009). *Appl. Phys. Lett. 95 231102*.

[10] Llopis, F., & Tobías, I. (2005). *Sol. Energy Mater. Sol. Cells, 87 481*.

**Acknowledgements**

76 Solar Cells - Research and Application Perspectives


## **Polycrystalline Cu(InGa)Se2/CdS Thin Film Solar Cells Made by New Precursors**

Alessio Bosio, Daniele Menossi, Alessandro Romeo and Nicola Romeo

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/51684

## **1. Introduction**

In the last five years photovoltaic modules production continued to be one of the rapidly growing industrial sectors, with an increase well in excess of 40% per year. This growth is driven not only by the progress in materials and technology, but also by incentives to sup‐ port the market in an increasing number of countries all over the world. Besides, the in‐ crease in the price of fossil fuels in 2008, highlighted the necessity to diversify provisioning for the sake of energy security and to emphasize the benefits of local renewable energy sour‐ ces such as solar energy. The high growth was achieved by an increase in production capaci‐ ty based on the technology of crystalline silicon, but in recent years, despite the already very high industrial growth rates, thin film photovoltaics has grown at an increasingly fast pace and its market share has increased from 6% in 2006 to over 12% in 2010. However, the ma‐ jority of photovoltaic modules installed today are produced by the well-established technol‐ ogy of monocrystalline and polycrystalline silicon, which is very close to the technology used for the creation of electronic chips. The high temperatures involved, the necessity to work in ultra-high vacuum and the complex cutting and assembly of silicon "wafers", make the technology inherently complicated and expensive. In spite of everything, silicon is still dominating the photovoltaic market with 90% of sales. Other photovoltaic devices based on silicon are produced in the form of "thin films" or in silicon ribbons; these devices are still in the experimental stage. Amorphous silicon is a technology that has been on the market for decades and it is by now clear that it does not keep the promises of change and develop‐ ment that were pledged when it was initially launched.

© 2013 Bosio et al.; licensee InTech. This is an open access article distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2013 Bosio et al.; licensee InTech. This is a paper distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Without resorting to sophisticated photovoltaic devices such as multi-junction solar cells, where the cost of production is high, thin film silicon modules were generally poor in con‐ version efficiency and demonstrated low stability. On the other hand, silicon is not an ap‐ propriate material for implementation as a thin film, both for the difficulties of processing (necessity of high temperatures) and the inherent characteristics of the semiconductor which, being an "indirect gap" material has a low absorption coefficient in the visible radia‐ tion region. Because of this, silicon must either be deposited in thick layers or it is necessary to use complex light trapping techniques. Beyond the use of silicon, thin film technology has the advantage to provide large-scale productions, in which the panel is the final stage of inline processes and not the assembly of smaller cells, as in the case of crystalline silicon or polysilicon wafer-based modules. The highest rates of production (in terms of square meters of modules per minute) have assumed since the '70s, that, in the future, in order to compete with traditional energy sources, there will be just thin film modules. However the effective start of industrial production of thin film modules was delayed until 2000 due to problems with the reproducibility of the results, the stability over time and scalability of the layer dep‐ osition on large areas. Overcoming these problems, the photovoltaic modules that use CuIn‐ GaSe2 (CIGS) and CdTe thin film technology are already being produced with a high quality and conversion efficiency (12-14%), with expected values up to 15% for the near future. The cell interconnection integrated into large area modules (0.6 x 1.2 m2 ), with very limited use of raw materials, can minimize the production cost, so that the thin film modules will soon be able to compete with conventional modules based on the silicon wafer.

In addition to lowering the cost/m2 of the cell area, thin film technology offers the possibil‐ ity to produce devices on flexible substrates. This extends the opportunity to installing modules by adapting them to the shape of the surface thereby achieving complete architec‐ tural integration.

Moreover, Cu(InGa)Se2polycrystalline thin film modules have successfully passed the longterm tests in outdoor conditions, demonstrating a very good stability over time.

Beyond the potential benefits as sociated with terrestrial applications we must also consider that the Cu(InGa)Se2 showed good resistance to ionizing radiation, much more if compared with crystalline silicon cells; furthermore, the cells can also be made on very light weight flexible substrates. For these reasons, this material is very promising for space applications. From this point of view, Cu(InGa)Se2 is one of the most promising materials used in thinfilm technology; not only for the reasons mentioned above, but perhaps more importantly, because it has reached very high efficiencies comparable to that obtained, up to now, with the best Silicon solar cells, at both cell and module level.

The highest Cu(InGa)Se2 solar cell efficiency of 20,3% with 0,5 cm2 total area was gained in 2010 by Jackson et al. [1] from Zentrum fuer Sonnenenergie of Wasserstoff-und-Forschung Baden-Wuerttemberg (ZSW), Germany. In addition, many companies have made modules with efficiencies above 12% up to the fantastic world record of 17,8 % obtained with 30*x*30cm2 modules by the "Solar Frontier" research group from Showa Shell Sekiyu KK (Ja‐ pan), which exceeds the previous record of 17,4% achieved by the Q-Cells subsidiary com‐ pany, Solibro Gmb H.

**Figure 1.** Schematic structure of a Cu(In,Ga)Se2-based solar cell.

Without resorting to sophisticated photovoltaic devices such as multi-junction solar cells, where the cost of production is high, thin film silicon modules were generally poor in con‐ version efficiency and demonstrated low stability. On the other hand, silicon is not an ap‐ propriate material for implementation as a thin film, both for the difficulties of processing (necessity of high temperatures) and the inherent characteristics of the semiconductor which, being an "indirect gap" material has a low absorption coefficient in the visible radia‐ tion region. Because of this, silicon must either be deposited in thick layers or it is necessary to use complex light trapping techniques. Beyond the use of silicon, thin film technology has the advantage to provide large-scale productions, in which the panel is the final stage of inline processes and not the assembly of smaller cells, as in the case of crystalline silicon or polysilicon wafer-based modules. The highest rates of production (in terms of square meters of modules per minute) have assumed since the '70s, that, in the future, in order to compete with traditional energy sources, there will be just thin film modules. However the effective start of industrial production of thin film modules was delayed until 2000 due to problems with the reproducibility of the results, the stability over time and scalability of the layer dep‐ osition on large areas. Overcoming these problems, the photovoltaic modules that use CuIn‐ GaSe2 (CIGS) and CdTe thin film technology are already being produced with a high quality and conversion efficiency (12-14%), with expected values up to 15% for the near future. The

of raw materials, can minimize the production cost, so that the thin film modules will soon

In addition to lowering the cost/m2 of the cell area, thin film technology offers the possibil‐ ity to produce devices on flexible substrates. This extends the opportunity to installing modules by adapting them to the shape of the surface thereby achieving complete architec‐

Moreover, Cu(InGa)Se2polycrystalline thin film modules have successfully passed the long-

Beyond the potential benefits as sociated with terrestrial applications we must also consider that the Cu(InGa)Se2 showed good resistance to ionizing radiation, much more if compared with crystalline silicon cells; furthermore, the cells can also be made on very light weight flexible substrates. For these reasons, this material is very promising for space applications. From this point of view, Cu(InGa)Se2 is one of the most promising materials used in thinfilm technology; not only for the reasons mentioned above, but perhaps more importantly, because it has reached very high efficiencies comparable to that obtained, up to now, with

2010 by Jackson et al. [1] from Zentrum fuer Sonnenenergie of Wasserstoff-und-Forschung Baden-Wuerttemberg (ZSW), Germany. In addition, many companies have made modules with efficiencies above 12% up to the fantastic world record of 17,8 % obtained with

pan), which exceeds the previous record of 17,4% achieved by the Q-Cells subsidiary com‐

modules by the "Solar Frontier" research group from Showa Shell Sekiyu KK (Ja‐

), with very limited use

total area was gained in

cell interconnection integrated into large area modules (0.6 x 1.2 m2

the best Silicon solar cells, at both cell and module level.

The highest Cu(InGa)Se2 solar cell efficiency of 20,3% with 0,5 cm2

tural integration.

80 Solar Cells - Research and Application Perspectives

30*x*30cm2

pany, Solibro Gmb H.

be able to compete with conventional modules based on the silicon wafer.

term tests in outdoor conditions, demonstrating a very good stability over time.

As we can see in figure 1,the CuInGaSe2/CdS thin film solar cell consists of 6 layers; this im‐ plies that in the overall structure there are at least 7 interfaces.

This made very complicated to understand the behavior of the final device and several research groups have tried to explain the properties of the cell by studying these interfa‐ ces in detail. On the other hand, when two different materials are put in contact there is an inter-diffusion of chemical elements from the one to the other and a sub sequent forma‐ tion ofa new thin layer between the two. This new layer is known as *hetero-interface*. The most important hetero-interface is the metallurgical hetero-junction between Cu(InGa)Se2 and CdS, but all the other interfaces have also an important role in the final performance of the cell.

Despite all efforts aimed to understand the behavior of the interfaces, Cu(InGa)Se2/CdS het‐ ero-junction still exhibits quite a few open problems and it is therefore subject to a margin of uncertainty in its progress. For this reason more detailed studies are needed to reach a com‐ plete understanding of all the phenomena regarding this remarkable device.

In this chapter we will describe the current state and the degree of understanding of the Cu(InGa)Se2 solar cells construction technology. In particular, after presenting a brief history of this device we will discuss the material and consider both the cells and the modules; after that we will focus particularly on the manufacturing techniques which have led to high-effi‐ ciency devices (cells and modules) and consequently, the different problems in herent to this material with particular attention to the scalability at an industrial level of the production process. Then we arrive at conclusions also talking about future perspectives.

## **2. A brief history**

The history of CuInSe2 begins with the research carried outin the Bell Telephone laborato‐ ries in the early 70'seven thoughits synthesis and characterization have already been stud‐ ied by Hahn in 1953 [2]. Along with new ternary chalcopyrite materials, it was also characterized by other groups [3]. The Bell Labs had grown crystals of a wide selection of these materials reporting their structural and electro-optics properties [4, 5]. In that period, a solar cell with an efficiency of 12% based on CdS evaporated onto a *p*-type CuInSe2 single crystal was realized [6]. In 1977, depositing by flash-evaporation a CdS thin film onto a single crystal of *p*-type CuGaSe2, a solar cell that exhibited an energy conver‐ sion efficiency of up to 7 % was realized [7].

CuInSe2 is a semiconducting compound of the I-III-VI2 family with a direct band gap of 1,05 eV. Its chalcopyrite structure makes a good match to wurtzite CdS with only 1,2% lattice mis match. This explains the good efficiency for the first time obtained with CuInSe2 single crys‐ tal and put in evidence that CuInSe2/CdS was the sixth system, along with junctions based on Si, GaAs, CdTe, InP, and CuxS that showed energy conversion efficiency up to 10%. Be‐ sides, CuInSe2 is a direct band gap semiconductor, which minimizes the requirements for minority carrier diffusion length, and exhibits the highest absorption coefficient (3*x*105 cm-1) in the visible region of the solar spectrum. These considerations make CuInSe2 the best-suit‐ ed material for the fabrication of an all polycrystalline thin film solar cell.

There has been relatively little effort devoted to devices realized on a CuInSe2 single crystal apart from this first work, because of the difficulty in growing high-quality crystals. But, the aforementioned properties of CuInSe2 channeled all the attention to thin-film solar cells be‐ cause of their intrinsic advantages. The first thin-film CuInSe2/CdS solar cell was fabricated by Kazmerski et al. in 1976 [8] by using films deposited by evaporation of CuInSe2 powder in excess of Se vapor. This solar cell showed an efficiency of about 4-5%.

We had to wait until 1981 when, in the Boeing laboratories, the first high-efficiency all thin film solar cell based on the system *n*-ZnCdS/*p*-CuInSe2 was realized with a conversion effi‐ ciency of about 9,4% and in 1985 they reached the efficiency of 11.4% [9].

Since the early 80's, ARCO Solar and Boeing have tackled the difficult issues involved with industrial production such as through put and yield. These efforts have led to many advan‐ cesin the technology of CuInSe2 solar cells.

The two groups have characterized their R&D approaches in different ways. The diversity of the two approaches consists basically in the CuInSe2(CISe) deposition methods, while the architecture of the device remains essentially the same.

The Boeing method includes the co-evaporation from separate sources of the single elements for CISe deposition. These films were deposited on alkali-free glass or ceramic covered by a thin layer of Mo, which acts as a positive electrode. The devices were finally completed by evaporating, on top of the CISe film, two layers of CdS (or ZnCdS), the first one was an in‐ trinsic layer and the second one heavily doped with indium in order to ensure a best photo current collection.

The two methods, introduced by Boeing and ARCO Solar, still remain the most common techniques for producing high efficiency cells and modules. Boeing was focused on co-evap‐ oration of individual elements from separate crucibles while ARCO Solar was more confi‐ dent in a two-stage process in which a low-temperature deposition of Cuand In was followed by a heat treatmentat high temperature in H2Se ambient.

With the "Boeing" basic structure, in 1996 [10] was reached the fantastic result for an all thin film solar cell: a conversion efficiency of 17,7%! This improvement was obtained by using CuGaXIn1-XSe2 as absorber layer. Effects of partial substitution of Ga for In appeared to be optimized for X=0,25. The band gap of the quaternary compound varies from 1,04 eV for X=0 to 1,7 eV for X=1; this means that the substitution of Ga for In causes an increase in open-circuit voltage, but a decrease in short-circuit current and fill factor and only for X=0,25 does the system reach the right equilibrium. Adjusting the Ga concentration profile into the absorber layer, in order to enhance the collection of the photo-generated carriers, it was pos‐ sible to fabricate thin film solar cells based on the CdS/CuGaInSe2 system with an efficiency of 18,8% in 1999 [11], of 19,2% in 2003 [12] and of 20.3% in 2011 [1]. This last result is the highest value for energy conversion efficiency in an all thin film photovoltaic device.


**Table 1.** Representative CuInSe2 and CuGaSe2 based solar cells.

**2. A brief history**

82 Solar Cells - Research and Application Perspectives

sion efficiency of up to 7 % was realized [7].

cesin the technology of CuInSe2 solar cells.

current collection.

architecture of the device remains essentially the same.

The history of CuInSe2 begins with the research carried outin the Bell Telephone laborato‐ ries in the early 70'seven thoughits synthesis and characterization have already been stud‐ ied by Hahn in 1953 [2]. Along with new ternary chalcopyrite materials, it was also characterized by other groups [3]. The Bell Labs had grown crystals of a wide selection of these materials reporting their structural and electro-optics properties [4, 5]. In that period, a solar cell with an efficiency of 12% based on CdS evaporated onto a *p*-type CuInSe2 single crystal was realized [6]. In 1977, depositing by flash-evaporation a CdS thin film onto a single crystal of *p*-type CuGaSe2, a solar cell that exhibited an energy conver‐

CuInSe2 is a semiconducting compound of the I-III-VI2 family with a direct band gap of 1,05 eV. Its chalcopyrite structure makes a good match to wurtzite CdS with only 1,2% lattice mis match. This explains the good efficiency for the first time obtained with CuInSe2 single crys‐ tal and put in evidence that CuInSe2/CdS was the sixth system, along with junctions based on Si, GaAs, CdTe, InP, and CuxS that showed energy conversion efficiency up to 10%. Be‐ sides, CuInSe2 is a direct band gap semiconductor, which minimizes the requirements for minority carrier diffusion length, and exhibits the highest absorption coefficient (3*x*105

in the visible region of the solar spectrum. These considerations make CuInSe2 the best-suit‐

There has been relatively little effort devoted to devices realized on a CuInSe2 single crystal apart from this first work, because of the difficulty in growing high-quality crystals. But, the aforementioned properties of CuInSe2 channeled all the attention to thin-film solar cells be‐ cause of their intrinsic advantages. The first thin-film CuInSe2/CdS solar cell was fabricated by Kazmerski et al. in 1976 [8] by using films deposited by evaporation of CuInSe2 powder

We had to wait until 1981 when, in the Boeing laboratories, the first high-efficiency all thin film solar cell based on the system *n*-ZnCdS/*p*-CuInSe2 was realized with a conversion effi‐

Since the early 80's, ARCO Solar and Boeing have tackled the difficult issues involved with industrial production such as through put and yield. These efforts have led to many advan‐

The two groups have characterized their R&D approaches in different ways. The diversity of the two approaches consists basically in the CuInSe2(CISe) deposition methods, while the

The Boeing method includes the co-evaporation from separate sources of the single elements for CISe deposition. These films were deposited on alkali-free glass or ceramic covered by a thin layer of Mo, which acts as a positive electrode. The devices were finally completed by evaporating, on top of the CISe film, two layers of CdS (or ZnCdS), the first one was an in‐ trinsic layer and the second one heavily doped with indium in order to ensure a best photo

ed material for the fabrication of an all polycrystalline thin film solar cell.

in excess of Se vapor. This solar cell showed an efficiency of about 4-5%.

ciency of about 9,4% and in 1985 they reached the efficiency of 11.4% [9].

cm-1)

Let's summarize the key enhancements to the method that gave the more efficient cells (coevaporation-Boeing).


increasing the photocurrent having enlarged the spectral response in the blue region wavelengths [14].


Several companies around the world are coming to the market with Cu(In,Ga)Se2-based modules. The more advanced features are briefly shown in Table 2, where one can distin‐ guish the two processes described above. One is the typical co-evaporation method of Shell Solar Industries, formerly Arco Solar and then Siemens Solar in California, Würth Solar and Solibro in Germany and Matsushita in Japan, with which these companies have announced modules efficiencies of around 12-13%. The selenization of metallic precursors in H2Se ambi‐ entis instead the technology used by Showa Shell that has announced module efficiency in excess of 14%.


**Table 2.** Some manufacturers of CuInGaSe2 thin-film photovoltaic modules, and their current market condition.

In addition a Phoenix Solar Holdings Corp subsidiary company, EPV Solar Corp, formerly Energy Photovoltaics, Inc. is using its own in-line evaporation process, International Solar Electric Technology (ISET) is developing a particle-based precursor for selenization, while Global Solar Energy (GSE) and MiaSolè are pursuing a process for roll-to-roll co-evapora‐ tion onto a flexible substrate with final efficiencies of around 15.5%.

Despite these encouraging results and the efforts made to develop the manufacturing proc‐ esses, there remains a large difference in efficiency between the laboratory-scale solar cells and mini-modules, and the best modules on the market. In part, this is due to the need to develop innovative equipment for large-area, high-throughput deposition required for man‐ ufacturing thin-film photovoltaic's. Presumably, this is because the material has only been studied for use in photovoltaic applications, and many of the advances in scientific under‐ standing of materials and technologies have been purely empirical.

However, in recent years, many enhancements in the scientific knowledge of the materi‐ als and device have been made and this has also led to evident improvements in manufac‐ turing technology.

## **3. The history of CuInSe2 and Cu(In,Ga)se2 at the ThiFiLab**

increasing the photocurrent having enlarged the spectral response in the blue region

**3.** The absorber energy gap was increased from 1.02 eV for CuInSe2 to 1.1–1.2 eV for Cu(In,Ga)Se2 by the partial substitution of In with Ga, leading to an important increase

**4.** Innovative absorber deposition processes were developed to obtain energy gap gradi‐

Several companies around the world are coming to the market with Cu(In,Ga)Se2-based modules. The more advanced features are briefly shown in Table 2, where one can distin‐ guish the two processes described above. One is the typical co-evaporation method of Shell Solar Industries, formerly Arco Solar and then Siemens Solar in California, Würth Solar and Solibro in Germany and Matsushita in Japan, with which these companies have announced modules efficiencies of around 12-13%. The selenization of metallic precursors in H2Se ambi‐ entis instead the technology used by Showa Shell that has announced module efficiency in

CIGS

1" wide metal sheet 0.6 x 1.2 0.66 x 1.61

Wuerth Solar, Ger. 14.8 (2007) 0.6 x 1.2 <13/11.7 Yes

Showa Shell, Japan 20 (2007) 0.6 x 1.2 14.2/11.8 Yes Honda Soltec, Japan 27 (2007/2008) 0.8 x 1.3 13/10 Yes Sulfur Cell, Ger. 5 (2007/2008) 0.65 x 1.25 8.2/7 Yes AVANCIS, Ger. 20 (from 2008) 0.65 x 1.6 13.1/12.2 Yes

**Table 2.** Some manufacturers of CuInGaSe2 thin-film photovoltaic modules, and their current market condition.

tion onto a flexible substrate with final efficiencies of around 15.5%.

In addition a Phoenix Solar Holdings Corp subsidiary company, EPV Solar Corp, formerly Energy Photovoltaics, Inc. is using its own in-line evaporation process, International Solar Electric Technology (ISET) is developing a particle-based precursor for selenization, while Global Solar Energy (GSE) and MiaSolè are pursuing a process for roll-to-roll co-evapora‐

**Glass-size (m x m) Efficiency %**

30 (2008) 0.5 x 1.2 --/9.4 No

25-30 (2009) 0.65 x 1.2 12 Yes

**max./med.**

10/8 10 13

**On the market**

Yes Yes Yes

ents improving the photovoltage and current collection [16, 17].

wavelengths [14].

84 Solar Cells - Research and Application Perspectives

in efficiency [15].

excess of 14%.

Bosch Solar CISTech (Johanna Solar, Ger.)

Global Solar, USA ISET, USA MiaSolè, USA

Hanergy China (Solibro, Ger.)

**Producers Prodn. capacity**

**MW/Year (since)**

4.2 (2006) Pilot plant 150 (2012) The Thin Film Laboratory (ThiFiLab) of the Physics Department at the University of Parma-Italy, started to work on CuInSe2 in 1986 with the aim of achieving high quality films suita‐ ble to realize an all thin film solar cell made entirely using only the sputtering technique [18-22]. The first approach was the sputtering deposition of CuInSe2 starting directly from stoichiometric targets. At this stage, great care was placed on the preparation of the sub‐ strates. As substrates 1 inch square Corning glasses 7059 were used. Glasses were covered with a 4 μm thick layer of Al deposited by sputtering at a substrate temperature of 350°C. The Al layer obtained was crystalline with an average grain size of 100 μm and with the (111) plane uniquely oriented along the glass surface. Since Al is not a good ohmic contact for *p*-type CuInSe2, the Al film was covered with either a 0.2 μm thick layer of Mo or a 0.5 μm layer of Au; both these films were deposited by sputtering. Au was deposited on Al at a substrate temperature higher than 500°C. In this way, Au reacts with Al forming the metallic compound AuAl2. This compound exhibits the cubic structure with a lattice constant of about 6 Å which is closer, in respect to any other metal, (Mo has a lattice constant of 3.15 Å) to the lattice constant of CuInSe2 (5.78 Å). Both Mo and AuAl2 thin films, prepared as previ‐ ously described, were used as substrates for the sputtering deposition of CuInSe2.

The RF magnetron sputtering system was provided with three targets like Mo, CuInSe2 and CdS and a rotatable heating-etching station whose temperature could be controlled up to 600°C. Mo and CdS targets were supplied by commercial suppliers while the CuInSe2 target was prepared in the ThiFiLabby synthesis in a high pressure furnace. In this case the liquid encapsulation technique was used, with B2O3 acting as the encapsulant. The crucibles had a flat bottom whose diameter was equal to that of the target holder in the sputtering system. In order to avoid the reaction between In and Se at a temperature lower than the B2O3 melt‐ ing point (450°C) a pre-reacted Cu-In alloy instead of the individual elements was used. The reaction of components and melting were carried out under a 50 atm N2 pressure. Several CuInSe2 targets with different Cu/In ratios and one with an excess of Se with respect to the metals have been investigated. The substrate temperature was varied between 300 and 550°C. It was found out that stoichiometry, crystallinity, carrier type and resistivity of CuInSe2 films depended on several deposition parameters, like the Ar pressure, R.F. power, bias applied to the substrate, substrate temperature and type of substrate, whether glass, Mo or AuAl2 (see figures 2-3).

Among these, the substrate temperature and the bias applied to the substrate seemed to be the most effective. Stoichiometric films with good crystallinity were generally grown at a substrate temperature of around 450°C, with the substrate bias kept at a minimum value (self-biased substrate). CuInSe2/CdS solar cells were fabricated as depicted in figure 4.

Mo, CuInSe2 and undoped CdS films were deposited in sequence in the same sputtering chamber without interrupting the vacuum [23, 24]. Mo and CuInSe2 films were deposited at the same substrate temperature, between 400 and 500°C. In some cases, CuInSe2 films were deposited on AuAl2 covered substrates. The undoped CdS film was deposited at 200°C sub‐ strate temperature with an H2 partial pressure in the sputtering chamber of about 1.5xl0-4 mbar, that corresponds to 3% of the Ar+H2 total pressure. These films exhibited resistivity in the range of 10-1 - 1 Ωcm and were highly transparent in the wavelength region between 1.6 μm and the absorption edge of CdS (0.52 μm).

After that, a mecanical scribing was done in order to obtain an active cell area of about 0.2 cm2 . The photovoltaic efficiency of such a cell was in the range of 5-8%. To achieve high effi‐ ciency cells it was necessary to use two CuInSe2 layers, the first one Cu-rich and the second In-rich [25-27]. The low efficiency obtained with this technique was essentially due to the segregation of binary phases, like Cu2Se in the first layer.

In order to overcome the segregation of binary phases some alternative techniques were studied; among them the selenization of stacked elemental layers seemed to be one of the most effective since the technique is very promising for large-scale application.

**Figure 2.** a) Resistivity of *p*-type CuInSe2 films as a function of the substrate temperature. The Cu/In ratio in the target was 0.8 and no bias was applied to the substrate. (b) Resistivity of *p*-type CuInSe2 films as a function of the Cu/In ratio in the target. The substrate temperature is 100°C and no bias was applied to the substrate. Redrawn from "Proceed‐ *ings of the 18 th IEEE Photovoltaic Specialists Conf., 1985, Las Vegas, Nevada, October 21-25. p. 1388-1392".*

**Figure 3.** a) Resistivity of *p*-type CuInSe2 films as a function of sputtering power. The Cu/In ratio in the target was 0.8 and the substrate temperature was 200°C. No bias was applied to the substrate. (b) Resistivity of both *p*- and *n*-type CuInSe2 films as a function of the negative bias applied to the substrate. The Cu/In ratio in the target was 1 and the substrate temperature was kept at 250°C. Redrawn from *"Proceedings of the 18 th IEEE Photovoltaic Specialists Conf., 1985, Las Vegas, Nevada, October 21-25. p. 1388-1392".*

**Figure 4.** Structure of the all-sputtered thin film CuInSe2 solar cell.

550°C. It was found out that stoichiometry, crystallinity, carrier type and resistivity of CuInSe2 films depended on several deposition parameters, like the Ar pressure, R.F. power, bias applied to the substrate, substrate temperature and type of substrate, whether glass, Mo

Among these, the substrate temperature and the bias applied to the substrate seemed to be the most effective. Stoichiometric films with good crystallinity were generally grown at a substrate temperature of around 450°C, with the substrate bias kept at a minimum value

Mo, CuInSe2 and undoped CdS films were deposited in sequence in the same sputtering chamber without interrupting the vacuum [23, 24]. Mo and CuInSe2 films were deposited at the same substrate temperature, between 400 and 500°C. In some cases, CuInSe2 films were deposited on AuAl2 covered substrates. The undoped CdS film was deposited at 200°C sub‐ strate temperature with an H2 partial pressure in the sputtering chamber of about 1.5xl0-4 mbar, that corresponds to 3% of the Ar+H2 total pressure. These films exhibited resistivity in the range of 10-1 - 1 Ωcm and were highly transparent in the wavelength region between 1.6

After that, a mecanical scribing was done in order to obtain an active cell area of about 0.2

In order to overcome the segregation of binary phases some alternative techniques were studied; among them the selenization of stacked elemental layers seemed to be one of the

**Figure 2.** a) Resistivity of *p*-type CuInSe2 films as a function of the substrate temperature. The Cu/In ratio in the target was 0.8 and no bias was applied to the substrate. (b) Resistivity of *p*-type CuInSe2 films as a function of the Cu/In ratio in the target. The substrate temperature is 100°C and no bias was applied to the substrate. Redrawn from "Proceed‐

*ings of the 18 th IEEE Photovoltaic Specialists Conf., 1985, Las Vegas, Nevada, October 21-25. p. 1388-1392".*

most effective since the technique is very promising for large-scale application.

. The photovoltaic efficiency of such a cell was in the range of 5-8%. To achieve high effi‐ ciency cells it was necessary to use two CuInSe2 layers, the first one Cu-rich and the second In-rich [25-27]. The low efficiency obtained with this technique was essentially due to the

(self-biased substrate). CuInSe2/CdS solar cells were fabricated as depicted in figure 4.

or AuAl2 (see figures 2-3).

86 Solar Cells - Research and Application Perspectives

cm2

μm and the absorption edge of CdS (0.52 μm).

segregation of binary phases, like Cu2Se in the first layer.

By using the same substrates described above, a 2500 Å thick Cu film and a 5600Å thick In film were deposited by sputtering on top of the Mo layer. The deposition of Cu and In ele‐ mental layers was done at several temperatures, ranging from 50 to 220°C. After the deposi‐ tion of the elemental layers, the sample was set inside an evaporation chamber for the selenization. The Se-vapor was obtained from a graphite boat kept 5 cm apart from the sub‐ strate holder, at a substrate temperature of about 300°C. The Se-deposition rate, measured by a quartz crystal monitor, was on the order of 0.5-0.8 μm/sec. The substrate was indirectly heated up using the Joule-effect with a Mo thin sheet. In this way a substrate temperature of 400-500°C could be reached in less than one minute. One of the problems encountered with the selenization of elemental layers was the sticking of the formed CuInSe2 film. If the In lay‐ er was deposited on top of Cu layer at a substrate temperature below 100°C the sticking was very poor. The best adhesion was obtained when the In layer was deposited at a substrate temperature close to its melting point. It seems that at this temperature Cu and In mix com‐ pletely [28]. After the selenization process, the CuInSe2 film was polycrystalline with a grain size larger than 1 micron and it was very well oriented along the (112) direction. The CuInSe2/CdS solar cells were made following the method described above. The best cells prepared with this technique gave the following photovoltaic parameters: Voc = 380 mV, Jsc = 39 mA/cm2 , f.f. = 0.61, η = 9.1%

The conclusion of this work was that the selenization of elemental Cu-In layers seemed to be a very good technique to make CuInSe2 films for the following three reasons:


It was evident that in order to improve the performance of the cells it was necessary to add Ga to the precursor. In this way a CuInxGa1-xSe2 film as the absorber material is obtained. Since the sputtering technique didn't allow Ga to be deposited directly, the deposition of the precursor was carried out by high vacuum evaporation assisted by an electron beam gun (E.B.G.).The Cu, Ga and In elemental layers were deposited by E.B.G. evaporation from a single rotatable crucible, in order to have a good uniformity on the substrate. First of all, 370 nm of In and 240 nm of Cu were deposited in sequence on top of a Mo covered glass sub‐ strate, which was kept at 120°C. Then the In-Cu bilayer was annealed at 290°C for 30-60 min in order to form the Cu11In9 phase. After the annealing, 120 nm of Ga were deposited at 240°C substrate temperature. At this temperature Ga reacts with the Cu11In9 phase forming the CuGa2 phase and freeing a small amount of In. A substrate temperature of 240°C was chosen for the deposition of Ga since the CuGa2 phase melts at 254°C. It had been confirmed by X-ray diffraction that the Cu-Ga-In film was made up of the Cu11In9 phase, the CuGa2 phase and a small amount of In. This method for preparing the Cu-Ga-In elemental layers gives the maximum possible intermixing between the layers, without having melt droplets in the film. Since the addition of Ga, being a low-melting element, increases the amount of liquid present in the film at a given temperature, a new selenization process was implement‐ ed. A peculiarity of this process was the oscillation of the substrate temperature at the first stage of CuGaInSe2 growth. The temperature oscillation seemed to be very effective in con‐ taining the surface tension of the liquid. At the beginning, the substrate temperature was raised up to 280°C in about 30 s and decreased to 200°C in about 3-4 min. During this first oscillation a small part of the film grew on the surface. This covering seemed to be sufficient to contain the surface tension during the second oscillation up to 450°C, which was done with a rise time of about 1 min and a cooling down time of about 6 min. During the second oscillation part of the film grew further covering the film with a solid one. Hence allowing a third temperature oscillation and so on. It has been seen that the process depends on both the oscillation rise time, which should not be larger than 2 min for reaching a maximum temperature of 550°C and the cooling down time from 550°C to 200°C which should be be‐ tween 6 to 9 min.

400-500°C could be reached in less than one minute. One of the problems encountered with the selenization of elemental layers was the sticking of the formed CuInSe2 film. If the In lay‐ er was deposited on top of Cu layer at a substrate temperature below 100°C the sticking was very poor. The best adhesion was obtained when the In layer was deposited at a substrate temperature close to its melting point. It seems that at this temperature Cu and In mix com‐ pletely [28]. After the selenization process, the CuInSe2 film was polycrystalline with a grain size larger than 1 micron and it was very well oriented along the (112) direction. The CuInSe2/CdS solar cells were made following the method described above. The best cells prepared with this technique gave the following photovoltaic parameters: Voc = 380 mV, Jsc =

The conclusion of this work was that the selenization of elemental Cu-In layers seemed to be

**b.** it allows the formation of well-crystallized and single phase CuInSe2 layers over large

It was evident that in order to improve the performance of the cells it was necessary to add Ga to the precursor. In this way a CuInxGa1-xSe2 film as the absorber material is obtained. Since the sputtering technique didn't allow Ga to be deposited directly, the deposition of the precursor was carried out by high vacuum evaporation assisted by an electron beam gun (E.B.G.).The Cu, Ga and In elemental layers were deposited by E.B.G. evaporation from a single rotatable crucible, in order to have a good uniformity on the substrate. First of all, 370 nm of In and 240 nm of Cu were deposited in sequence on top of a Mo covered glass sub‐ strate, which was kept at 120°C. Then the In-Cu bilayer was annealed at 290°C for 30-60 min in order to form the Cu11In9 phase. After the annealing, 120 nm of Ga were deposited at 240°C substrate temperature. At this temperature Ga reacts with the Cu11In9 phase forming the CuGa2 phase and freeing a small amount of In. A substrate temperature of 240°C was chosen for the deposition of Ga since the CuGa2 phase melts at 254°C. It had been confirmed by X-ray diffraction that the Cu-Ga-In film was made up of the Cu11In9 phase, the CuGa2 phase and a small amount of In. This method for preparing the Cu-Ga-In elemental layers gives the maximum possible intermixing between the layers, without having melt droplets in the film. Since the addition of Ga, being a low-melting element, increases the amount of liquid present in the film at a given temperature, a new selenization process was implement‐ ed. A peculiarity of this process was the oscillation of the substrate temperature at the first stage of CuGaInSe2 growth. The temperature oscillation seemed to be very effective in con‐ taining the surface tension of the liquid. At the beginning, the substrate temperature was raised up to 280°C in about 30 s and decreased to 200°C in about 3-4 min. During this first oscillation a small part of the film grew on the surface. This covering seemed to be sufficient to contain the surface tension during the second oscillation up to 450°C, which was done with a rise time of about 1 min and a cooling down time of about 6 min. During the second oscillation part of the film grew further covering the film with a solid one. Hence allowing a

a very good technique to make CuInSe2 films for the following three reasons:

**c.** Se-vapor in the selenization process works as well as H2Se.

39 mA/cm2

areas.

, f.f. = 0.61, η = 9.1%

88 Solar Cells - Research and Application Perspectives

**a.** it is a scalable technique.

At this stage, the front-contact was also changed; 30-50 nm thick layer of intrinsic CdS at 220°C substrate temperature was deposited by sputtering on top of the CuGa0.3In0.7Se2 films. The cell was completed with a 1-2 μm thick film of In-doped CdS or Al-doped ZnO, deposit‐ ed by sputtering. On average, better results were obtained when soda lime glass was used as a substrate, indicating that the diffusion of Na atoms into the growing CuGa0.3In0.7Se2 film during the selenization step, should play some role. These cells shown a maximum efficien‐ cy of 12.4% with an open-circuit voltage of 450 mV, a short-circuit current density of 40 mA/cm2 and a fill factor of 0.69. No better efficiencies were achieved because this process doesn't allow a Ga concentration gradient to be obtained within the film thickness. Proba‐ bly, if the junction was made with a thin layer of CdS obtained by chemical bath deposition, a larger efficiency could be achieved [29]. For this reason new type of precursors were test‐ ed, which allowed the concentration profile of Ga to be easily designed[30].

As it was the custom at that time, a substrate of l inch2 , 4mm thick soda-lime glass was used. The substrate was mounted in a sputtering chamber, where 4 targets namely Mo, In2Se3, Cu and a Cu-Ga alloy containing 50% at. ofGa and 50% at. of Cu were installed [31, 32]. Com‐ mon commercial suppliers provided Mo, Cu and the Cu-Ga alloy, while the In2Se3 target was home-prepared, starting from the In and Se elements. In and Se elements were put in a suitable graphite container and brought to a temperature higher than the In2Se3 melting point (890°C) and then slowly cooled down to room temperature. This process was done in an oven where an inert gas such as N2 was introduced at a pressure higher than 10 atmos‐ pheres in order to avoid the In or Se evaporation. The target was polycrystalline and exhibit‐ ed a 99.9% density comparable with the bulk material.

This characteristic of the target is very important since, if a target prepared by "hot pressing" starting from In2Se3 powder was used one could have instabilities in the sputtering dis‐ charge produced by the high vapor pressure of Se.

Furthermore, if the cooling of the target is not very effective, some cracks at the edges of the target can form. On top of the soda-lime glass, 1 μm of Mo, 1.5 μm of In2Se3 and 0.2 μm of Cu were deposited in sequence at room temperature by sputtering. Mo and Cu were depos‐ ited by pulsed D.C. sputtering while, In2Se3 was deposited by R.F. sputtering.

The Mo-In2Se3-Cu stacked layers were brought to a temperature of 450°C where they were left to interact, respecting their phase diagram, for about 30 min. In this way a film of CuInSe2 was formed. As it can be seen in figure 5(a), this layer contains a residual of the Cu11In9 and In2Se3 phases.

**Figure 5.** a) CuInSe2 precursor X-Ray Diffraction spectrum. (b) Comparison between CuInSe2 (dashed- line) and CuIn‐ GaSe2 (solid-line) X-Ray Diffraction spectra.

In order to obtain the right CuInSe2 film without any secondary phases, the precursor had to be selenized at a substrate temperature of 520-530°C. The selenization was done in the same way described before. Similar technology was also reported by other authors [33] but InSe instead of In2Se3 was used and it was deposited by high vacuum evaporation (HVE).

Furthermore, the selenization was started from room temperature up to 500°C and then no precursor was formed before selenization. In order to make the Cu(In,Ga)Se2 film, the CuInSe2 layer, obtained as described before, was covered by 80 nm of the Cu-Ga alloy (50%-50%) deposited by D.C. sputtering and this system was again selenized. The Cu(In,Ga)Se2 film, formed in this way has a graded Ga content, that means it contains more Ga close to the surface than into the bulk. The Cu(In,Ga)Se2, analyzed by X-rays, exhibits the peaks of Cu(In,Ga)Se2 material as one can see from figure 5(b) where the diffraction peaks are shifted to higher angles with respect to CuInSe2.

**Figure 6.** a) CuInSe2 film surface morphology after selenization. (b) CuInGaSe2 film surface morphology after selenization.

Figure 6(a) shows the morphology of the CuInSe2 film while figure 6(b) displays the mor‐ phology of the Cu(In,Ga)Se2 film referred to the two-steps process presented in figure 7. In order to have a better control on the Ga concentration gradient of the CuInGaSe2 film, the ThiFiLab developed an alternative way to prepare the precursor (see figure 7).

Once the CuInSe2 precursor is obtained it is possible to follow two different pathways:


**Figure 5.** a) CuInSe2 precursor X-Ray Diffraction spectrum. (b) Comparison between CuInSe2 (dashed- line) and CuIn‐

In order to obtain the right CuInSe2 film without any secondary phases, the precursor had to be selenized at a substrate temperature of 520-530°C. The selenization was done in the same way described before. Similar technology was also reported by other authors [33] but InSe

Furthermore, the selenization was started from room temperature up to 500°C and then no precursor was formed before selenization. In order to make the Cu(In,Ga)Se2 film, the CuInSe2 layer, obtained as described before, was covered by 80 nm of the Cu-Ga alloy (50%-50%) deposited by D.C. sputtering and this system was again selenized. The Cu(In,Ga)Se2 film, formed in this way has a graded Ga content, that means it contains more Ga close to the surface than into the bulk. The Cu(In,Ga)Se2, analyzed by X-rays, exhibits the peaks of Cu(In,Ga)Se2 material as one can see from figure 5(b) where the diffraction peaks

**Figure 6.** a) CuInSe2 film surface morphology after selenization. (b) CuInGaSe2 film surface morphology after selenization.

Figure 6(a) shows the morphology of the CuInSe2 film while figure 6(b) displays the mor‐ phology of the Cu(In,Ga)Se2 film referred to the two-steps process presented in figure 7. In

instead of In2Se3 was used and it was deposited by high vacuum evaporation (HVE).

GaSe2 (solid-line) X-Ray Diffraction spectra.

90 Solar Cells - Research and Application Perspectives

are shifted to higher angles with respect to CuInSe2.

With this second method we have more freedom in varying the concentration profile of Ga in the final Cu(In,Ga)Se2 film since the profile depends on the ratio between the Ga2Se3 and CuInSe2 film thicknesses. In contrast, the first method is more convenient because it needs only one selenization step. In Table 3 the results of manufactured cells with both methods are summarized.

Great improvements have been made in finishing the cell, in fact Cu(In,Ga)Se2/CdS solar cells were prepared as described above, by depositing in sequence on top of Cu(In,Ga)Se2 layer, 60 nm of CdS(F) [34], 80 nm of pure ZnO and 1μm of ZnO doped with 2% atomic of Al.

**Figure 7.** Representation of the two options developed in the THIFILAB, University of Parma-Italy to obtain high quali‐ ty Cu(In,Ga)Se2 absorber film. (a) single-step process. (b) double-step process. For both options all the layers are depos‐ ited by sputtering and selenization in pure Se atmosphere is made.

If one uses the sputtering technique for the deposition of CdS films, employed as the win‐ dow material in Cu(In,Ga)Se2 based solar cells, great results are not obtained and efficiencies converge to values not higher than 13% -14%. This result is due to the fact that these devices exhibit a too high diode reverse saturation current. One possible explanation is that the grain boundaries in the CdS film are active and can channel the diode reverse current.


**Table 3.** Photovoltaic parameters of the solar cells fabricated by sputtering and selenization in pure Se atmosphere.

It is possible to get over this problem by introducing in the sputtering chamber, during the CdS deposition, Argon containing 3% of CHF3. This gas is decomposed and ionized in the sputtering discharge, freeing F ions which, being strongly electronegative, are directed to the substrate that is the positive electrode; here two different events can happen:

1- the presence of energetic F- ions near the substrate favors the formation of a fluorine com‐ pound such as CdF2 during the growth of the CdS film [34].

2- the F ions, accelerated by the electric field present in the discharge, hit the film surface during the deposition with sufficient energy to sputter back the more weakly bonded Cd or S atoms.

This effect leaves a CdS film with high optical quality and structural properties. We can see in figure 8 that the CdS films deposited in Argon+CHF3 have an energy gap greater than that of the films deposited in Argon alone. Better efficiencies (15%-18%) are routinely ob‐ tained if CdS(F) films are used. One can explain this fact by considering that CdS(F) may contain CdF2 probably segregated in the grain boundaries and this can be useful to passi‐ vate them.

As an alternative to ZnO(Al), 0.5 μm of ITO doped with Zr was used. All the layers were deposited by sputtering. While CdS(F) and pure ZnO were deposited by R.F., ZnO(Al) and ITO(Zr) were deposited by pulsed D.C. sputtering. ITO doped with Zr has been used since it was discovered that for dielectric oxides, permittivity can be increased by the addition of higher-permittivity oxides such as ZrO2 (or HfO2). It is also known that the permittivity of dielectric oxides increases rapidly with even small additions of a higher-permittivity constit‐ uent and this is especially true for high-frequency permittivity ε∞. The increase in ε∞ as Zr is added can shift λp, which corresponds to the plasma resonance wavelength, to a longer wavelength and make it possible to improve NIR transmission significantly without altering the material parameters, like carrier concentration or mobility[35].

If one uses the sputtering technique for the deposition of CdS films, employed as the win‐ dow material in Cu(In,Ga)Se2 based solar cells, great results are not obtained and efficiencies converge to values not higher than 13% -14%. This result is due to the fact that these devices exhibit a too high diode reverse saturation current. One possible explanation is that the grain boundaries in the CdS film are active and can channel the diode reverse current.

**Absorber Ga/In+Ga (%) Voc (mV) Jsc (mA/cm2) FF (%) Efficiency (%)** CuInSe2 0 410 39.8 70.20 11.45 Cu(In,Ga)Se2 20 531 34.8 71.60 13.23 Cu(In,Ga)Se2 20 578 35.2 71.08 14.46

**Table 3.** Photovoltaic parameters of the solar cells fabricated by sputtering and selenization in pure Se atmosphere.

It is possible to get over this problem by introducing in the sputtering chamber, during the CdS deposition, Argon containing 3% of CHF3. This gas is decomposed and ionized in the

1- the presence of energetic F- ions near the substrate favors the formation of a fluorine com‐

during the deposition with sufficient energy to sputter back the more weakly bonded Cd or

This effect leaves a CdS film with high optical quality and structural properties. We can see in figure 8 that the CdS films deposited in Argon+CHF3 have an energy gap greater than that of the films deposited in Argon alone. Better efficiencies (15%-18%) are routinely ob‐ tained if CdS(F) films are used. One can explain this fact by considering that CdS(F) may contain CdF2 probably segregated in the grain boundaries and this can be useful to passi‐

As an alternative to ZnO(Al), 0.5 μm of ITO doped with Zr was used. All the layers were deposited by sputtering. While CdS(F) and pure ZnO were deposited by R.F., ZnO(Al) and ITO(Zr) were deposited by pulsed D.C. sputtering. ITO doped with Zr has been used since it was discovered that for dielectric oxides, permittivity can be increased by the addition of higher-permittivity oxides such as ZrO2 (or HfO2). It is also known that the permittivity of dielectric oxides increases rapidly with even small additions of a higher-permittivity constit‐ uent and this is especially true for high-frequency permittivity ε∞. The increase in ε∞ as Zr is added can shift λp, which corresponds to the plasma resonance wavelength, to a longer wavelength and make it possible to improve NIR transmission significantly without altering

ions, accelerated by the electric field present in the discharge, hit the film surface

the substrate that is the positive electrode; here two different events can happen:

pound such as CdF2 during the growth of the CdS film [34].

the material parameters, like carrier concentration or mobility[35].

ions which, being strongly electronegative, are directed to

sputtering discharge, freeing F-

92 Solar Cells - Research and Application Perspectives

2- the F-

S atoms.

vate them.

**Figure 8.** Transmission spectrum of an 80 nm thick sputtered CdS film: (a) deposited in pure Argon and (b) deposited in Argon+CHF3. The shift of the absorption edge towards shorter wavelengths proves the beneficial effects of deposi‐ tion in the presence of CHF3.Redrawn from *"Proceedings of 3rdWorld Conference on Photovoltaic Energy Conversion*, 2003, Osaka, Japan, May18-21, Vol 1, 469 – 470.

In order to further improve the efficiency of the solar cells, the precursor was further modi‐ fied by replacing In2Se3and Ga2Se3 with InSe and GaSe targets. The change in the starting se‐ lenides was especially effective as it resulted in the complete mixing of the materials, which formed the precursor. In fact, at a given temperature, InSe and GaSe were more reactive with Cu than their counter parts In2Se3 and Ga2Se3. So we came to the last version of the fab‐ rication process for the CIGS-baseds olar cells developed so far at the ThiFiLab.

Thestarting systemwas the usualsputtering machinein which 4 targets namely Mo, InSe, GaSe and Cu were contained. While Mo and Cu were deposited by pulsed-D.C., InSe and GaSe were deposited by R.F. sputtering. The constructionof the solar cellbeganwith thede‐ position of a Mo bi-layer at room temperature. The first Mo layer was quite thin, 30 nm and was deposited with an Ar flow of 45 sccm, which corresponds to apressureof 4.5x 10-3mbar. The Ar flow was then decreased to 15 sccm, which corresponds to apressureof 1.5x 10-3mbar.

Approximately 500 nm of Mo was deposited. The deposition of a bi-layer of Mo was neces‐ sary in order to have good adhesion of the Mo film to the substrate [36].

In order to get a good precursor, it was discovered that the sequence of InSe, GaSe and Cu as well as the substrate temperature at which they were deposited wasvery important. While Mo films had to be deposited at room temperature for sticking purposes, both InSe and GaSe layers were deposited at a high temperature (400°C) in order to avoid a Se-excess growth. The thickness of the InSe film was commonly 1.5 μm and that of GaSe was 0.5 μm. Cu was also deposited at asubstrate temperature of 400°C and its thickness was around 350 nm. All the layers mix during the deposition and don't require further annealing. An X-ray analysis shows that the precursor exhibits a mix of InSe, GaSe and CuInSe2 phases (figure 9a), while, the selenized material shows two phases one which contains 30% of Ga and the other one which contains 60% of Ga(figure 9b).

**Figure 9.** Cu(In,Ga)Se2 precursor X-Ray Diffraction spectra (a) before and (b) after selenization.

**Figure 10.** Electron microscope image of precursor (a) before and (b) after selenization.

The precursor is then selenized in a vacuum chamber where pure Se is evaporated from a graphite crucible. Selenization lasts approximately 7 minutes. The first 5 minutes are used to bring the substrate from room temperature to 530°C and the last 2 minutes are spent by leaving the substrate at 530°C. The morphology of the precursor before and after seleniza‐ tion is shown in figure 10.

**Figure 11.** Preparation of the precursor. (a) Sequence of deposition for the InSe, GaSe and Cu layers; (b) the last new optimized sequence for the InSe, GaSe and Cu layers, which led to a high efficiency CIGS-based so‐ lar cell.

The cell was completed by depositing the CdS layer, the transparent electrode and the con‐ tact grid in the same way that was previously described.

Solar cells made in this way exhibit an efficiency of about 13% with an open circuit voltage (VOC) that is never higher than 500 mV, a high short circuit current density (JSC) on the order of 40 mA/cm2 and a f.f. of 0.62-0.65.

**Figure 9.** Cu(In,Ga)Se2 precursor X-Ray Diffraction spectra (a) before and (b) after selenization.

94 Solar Cells - Research and Application Perspectives

**Figure 10.** Electron microscope image of precursor (a) before and (b) after selenization.

The low open circuit voltage and f.f. are attributed to the fact that Ga, being less reactive than In, tends to diffuse to the bottom leaving the surface poor in Ga. This has been con‐ firmed by doing a depth profile, in which one can see that most of Ga is confined to the bot‐ tom. At this point, in order to have more Ga close to the surface, a final change in the sequence of the layers had been performed and is highlighted in figure 11.

A GaSe thin layer is put on top of Cu and another layer of GaSe is put under Cu on top of InSe. The new precursor has been selenized in the same way as the old precursor. A depth profile made on this new absorber shows that there is more Ga close to the sur‐ face and that it decreases starting from the surface and increases again going deep into the sample (figure 12). The characteristics of the solar cells made with the new absorber exhib‐ it a VOC close to 570 mV, a JSC of ~38 mA/cm2 and f.f. of 0.74 with an efficiency of 16.2% (figure 13).

**Figure 12.** Sims depth profile of a CIGS film obtained starting from the new stratigraphy of the InSe/GaGebased precursor. Note the trend of the Ga profile, which decreases starting from the surface and then re-in‐ crease approaching the back contact.

**Figure 13.** J-V characteristics of the Cu(In,Ga)Se2 based solar cell made with the new precursor.

The process described above was also used to prepare Cu(In,Ga)Se2/CdS solar cells directly on ceramic substrates (commercial tiles). This kind of ceramics is normally used in construc‐ tion of buildings aimed to energy savings through the implementation of ventilated walls. However, the ceramic tile was modified in order to adapt its surface to become a good sub‐ strate for PV purpose. For this reason the surface of the tile was vitrified with the applica‐ tion of a special enamel. By checking in detail the composition and the constituent elements of this layer it has been possible to make the ceramic very similar, from the physico-chemi‐ cal point of view, to the most common soda-lime glass. In fact, solar cells prepared on this ceramic substrate exhibit similar results to those obtained by using soda-lime glass as a sub‐ strate. (see figure 14 and table 4)

**Figure 14.** Scanning electron microscope photograph of the surface morphology of a CIGS film grown on an evolved ceramic tile.


**Table 4.** Photovoltaic parameters of CIGS-based solar cells fabricated by using the last two different precursors both on soda-lime glass and commercial ceramic tile substrates.\*the absorber layer is made starting with the In2Se3+Ga2Se3+Cu precursor and the cell is realized on SLG. # the absorber layer is made starting with the InSe+GaSe +Cu+GaSe precursor and the cell is realized on SLG.^the absorber layer is made starting with the InSe+GaSe+Cu+GaSe precursor and the cell is realized on commercial ceramic tile.

## **4. A future perspective: the cu2znsns4 system**

**Figure 12.** Sims depth profile of a CIGS film obtained starting from the new stratigraphy of the InSe/GaGebased precursor. Note the trend of the Ga profile, which decreases starting from the surface and then re-in‐

**Figure 13.** J-V characteristics of the Cu(In,Ga)Se2 based solar cell made with the new precursor.

The process described above was also used to prepare Cu(In,Ga)Se2/CdS solar cells directly on ceramic substrates (commercial tiles). This kind of ceramics is normally used in construc‐ tion of buildings aimed to energy savings through the implementation of ventilated walls. However, the ceramic tile was modified in order to adapt its surface to become a good sub‐ strate for PV purpose. For this reason the surface of the tile was vitrified with the applica‐ tion of a special enamel. By checking in detail the composition and the constituent elements of this layer it has been possible to make the ceramic very similar, from the physico-chemi‐

crease approaching the back contact.

96 Solar Cells - Research and Application Perspectives

Since it is necessary to develop non-Si-based solar cells due to a lack of highly pure Si sour‐ ces, the PV world is oriented to develop new type of solar cell material that uses simple process and low-cost easily scalable techniques. Thin film polycrystalline solar cells have shown a remarkable growth in terms of efficiency, stability and scalability. CdTe and CuIn‐ GaSe2 (CIGS) materials have demonstrated to be robust and reliable for delivering low cost solar electricity. Because of direct band gap, which assures a high absorption coefficient, these devices use a very limited amount of material, which strongly reduces the production costs. However some issues on this are raised in particular the presence of rare and toxic ele‐ ments that could impact on the perception of the population.

Indium and gallium for CIGS and tellurium for CdTe result to be rare materials that could limit the module production when we talk about Terawatts, if a proper recycling system is not set up [37].

Moreover, there is an old controversial discussion on the opportunity of using cadmium in photovoltaics; unfortunately this argument has been going above the scientific considera‐ tions and has been taken as an opportunity for crystalline silicon producer to limit the spreading of less expensive polycrystalline thin film devices [38].

As a matter of fact the introduction of cadmium in photovoltaic materials has been in the fabrication either of CdTe absorber material or of CdS buffer layer; in the first case there is mainly a problem or perception since CdTe is a secure and non toxic material, in the second case the amount of material is very little and encapsulated and the impact of cadmium in the overall environmental effect is minimal.

However, to overcome the environmental issues and most important to remove any doubt of massive production connected with material scarcity, in 1996 Katagiri et al. [39] have in‐ troduced a new device based on Cu2ZnSnS4 (CZTS) absorber layer; the structure was very similar to the CIGS standard structure having CdS as buffer layer, ZnO as front and Mo as back contact. The idea was not new since it has been introduced for the first time in 1988 from Ito and Nagazawa by making a heterojunction between CZTS and CTO (Cadmium Tin Oxide) [40]. The introduction of CZTS allows practically to substitute indium and gallium with zinc and tin which are pretty much available on the earth crust.

Furthermore, thin film photovoltaics is competing with other industries, which use similar elements such as In (flat panel) and Ga (optoelectronics). This issue together with the raw material availability mentioned above could affect the raw material prices. As already writ‐ ten the availability of In could limit the module production in the terawatt/year range [37].

The advantage of the kesterite materials (CZTS family) is that with similar material proper‐ ties, similar preparation methods and same device structure as chalcopyrite solar cells, they are a real alternative to In-containing absorbers.

Generally for kesterites two kind of absorbers are considered: Cu2ZnSnS4 and Cu2ZnSnSe4, however their structure is pretty different since in the first case we have a real kesterite structure (space group I4) and in the second case a stannite structure is observed (space group I\_42m).

So when it comes to a Cu2ZnSn(S,Se)4 absorber, which has also been employed, a polymorph structure has been observed. Polymorphism gives place to coexistence of different structure and it has been shown that stannite structure has lower binding energy and so lower stabili‐ ty. Stannite structure has lower band gaps and this could explain the lower open circuit vol‐ tages of these devices compared to the expected values. Band gap of the materials are fluctuating depending on the different deposition methods and crystallization procedure; however most of band gaps for Cu2ZnSnS4 are around 1.5eV [41-46] while for Cu2ZnSnSe4 they stay around 1 eV. The main reason for the fluctuation of the band gaps is the very like‐ ly presence of secondary phases. As a matter of fact CZTS materials tend to produce many secondary phases, investigation on the phase diagram of this material shows that only in a very small region is possible to make single phase CZTS while a large variety of different secondary phases can easily be formed, such as ZnS(e), CuS(e), CuSnS(e). However the best solar cells show a Zn-rich and Cu-poor composition, which tendentially provides a ZnS(e) secondary phase. While ZnS(e), having a wide band gap and low conductivity, would not affect the open-circuit voltage, it could be responsible for the high series resistance observed in all solar cells [47,48].

solar electricity. Because of direct band gap, which assures a high absorption coefficient, these devices use a very limited amount of material, which strongly reduces the production costs. However some issues on this are raised in particular the presence of rare and toxic ele‐

Indium and gallium for CIGS and tellurium for CdTe result to be rare materials that could limit the module production when we talk about Terawatts, if a proper recycling system is

Moreover, there is an old controversial discussion on the opportunity of using cadmium in photovoltaics; unfortunately this argument has been going above the scientific considera‐ tions and has been taken as an opportunity for crystalline silicon producer to limit the

As a matter of fact the introduction of cadmium in photovoltaic materials has been in the fabrication either of CdTe absorber material or of CdS buffer layer; in the first case there is mainly a problem or perception since CdTe is a secure and non toxic material, in the second case the amount of material is very little and encapsulated and the impact of cadmium in the

However, to overcome the environmental issues and most important to remove any doubt of massive production connected with material scarcity, in 1996 Katagiri et al. [39] have in‐ troduced a new device based on Cu2ZnSnS4 (CZTS) absorber layer; the structure was very similar to the CIGS standard structure having CdS as buffer layer, ZnO as front and Mo as back contact. The idea was not new since it has been introduced for the first time in 1988 from Ito and Nagazawa by making a heterojunction between CZTS and CTO (Cadmium Tin Oxide) [40]. The introduction of CZTS allows practically to substitute indium and gallium

Furthermore, thin film photovoltaics is competing with other industries, which use similar elements such as In (flat panel) and Ga (optoelectronics). This issue together with the raw material availability mentioned above could affect the raw material prices. As already writ‐ ten the availability of In could limit the module production in the terawatt/year range [37]. The advantage of the kesterite materials (CZTS family) is that with similar material proper‐ ties, similar preparation methods and same device structure as chalcopyrite solar cells, they

Generally for kesterites two kind of absorbers are considered: Cu2ZnSnS4 and Cu2ZnSnSe4, however their structure is pretty different since in the first case we have a real kesterite structure (space group I4) and in the second case a stannite structure is observed (space

So when it comes to a Cu2ZnSn(S,Se)4 absorber, which has also been employed, a polymorph structure has been observed. Polymorphism gives place to coexistence of different structure and it has been shown that stannite structure has lower binding energy and so lower stabili‐ ty. Stannite structure has lower band gaps and this could explain the lower open circuit vol‐ tages of these devices compared to the expected values. Band gap of the materials are

ments that could impact on the perception of the population.

spreading of less expensive polycrystalline thin film devices [38].

with zinc and tin which are pretty much available on the earth crust.

overall environmental effect is minimal.

98 Solar Cells - Research and Application Perspectives

are a real alternative to In-containing absorbers.

not set up [37].

group I\_42m).

The standard structure of the solar cells is very similar to the CIGS one, using ZnO doped with aluminum as front contact, CdS as buffer layer and Molibdenum as back contact. Friedlmeier et al. introduced this first design in 1997; they fabricated thin film solar cells by using a CZTS layer as the light absorber in contact with an n-CdS/ZnO window layer [47]. The best energy conversion efficiency produced by these cells was 2.3%. Later, Katagiri's group broke this record in 1999 producing a CZTS solar cell with 2.63% power conversion efficiency. In this cell, the CZTS film was deposited on a Mo coated soda lime glass (SLG) substrate [42]. By optimization of the sulfurization process, the efficiency of the solar cells was increased to 5.45% in 2003 [49], and then to 6.7% in 2008 [50].

Finally in 2010 Todorov reached 9.6% by inserting a selenization step [51], in 2011 the psycho‐ logical limit of 10% has been overcome by Mitzi et al. by a non vacuum deposition method [52].

CZTS has been prepared already by a wide variety of deposition methods with different re‐ sults, however what is surprisingly remarkable is that both vacuum and non vacuum techni‐ ques are very much in competition on the efficiency level. This is not the case for other polycrystalline thin films where generally the vacuum techniques are delivering higher per‐ formance. A large number of different vacuum deposition methods have been employed.

Vacuum-based fabrication techniques normally involve deposition of the constituent atoms of the CZTS compound on a substrate by sputtering or evaporation/coevaporation of the tar‐ get sources under a certain pressure and temperature. These techniques have the advantage of easily controlling the chemical composition and phase profile in the thin films and nor‐ mally have good reproducibility.

As mentioned before the first CZTS processing by Ito and Nakazawa was made by vacuum deposition, such as atomic beam sputtering in 1988. They were able to measure a band gap of 1.45 eV and the photovoltaic effect was proved by preparing an heterojunction with cad‐ mium tin oxide (CTO) and a voltage of 165 mV was obtained [40].

After a long time in 1996 Katagiri made a complete photovoltaic cell with the following structure: ZnO:Al/CdS/CZTS/Mo/SLG. The absorber layer was deposited by electron beam gun followed by sulfurization; this process gave a power conversion efficiency of 0.66%. In 1997 Friedlmeier et al. at the university of Stuttgart deposited the CZTS absorber layer by coevaporation of the single constituents (using a similar method as the CIGS fabrication process where the single elements or binary chalcogenides compounds are vacuum-deposit‐ ed onto a heated substrate) obtaining a 2.3% efficiency cell. They also report, for the first time, that the electrical resistivity of the films can be improved considerably by treating the films with a KCN solution; this is still an important step of the fabrication process. Later on, Katagiri improved his deposition process by substituting Zn with ZnS in the e-beam se‐ quence and improved the adhesion of the absorber to the molybdenum, increasing its effi‐ ciency up to 1.08%.

An approach based on one-step synthesis of CZTS thin films by simultaneous coevaporation of precursor sources Cu, Zn, Sn, S was reported by Tanaka et al. The grain size of the film depended on the substrate temperature. Larger grains were obtained at higher substrate temperature [52].

An integrated vacuum apparatus, which combined the RF sputtering technique with the sul‐ furization process was employed to reduce the effect of moisture from air ambient on the property of CZTS films. Target sources were based on ZnS, SnS and Cu. Both the quality and reproducibility of the CZTS films were significantly improved by this method and a 5.7% conversion efficiency with the CZTS-based thin film solar cells was achieved [53]. By eliminating the metal oxide impurity in the CZTS films using deionized water, the efficiency of the solar cells was further improved to 6.7% [49,54].

A lot of work has been recently presented by the IBM group; the research has been going on two different paths: vacuum and non vacuum deposition of CZTS.

For vacuum deposition a recent work from Wang et al. has been presented where the single elements are contemporarily evaporated on a SLG substrate by a rapid thermal deposition system. With an absorber thickness of only 1 micron and a few minutes annealing time the process is very promising for industrial application and has demonstrated a 6.8% conver‐ sion efficiency. High series resistance and a high charge recombination limit the efficiency [55].

Generally, if the composition of the precursor films is copper rich, it results in the forma‐ tion of a Cu2−x S secondary phase, which can be removed by a post treating with KCN aqueous solution.

In metallic layer deposition it has been shown that the sequence of the different elements plays an important role in the crystallization of the absorber as shown with controversial re‐ sults by Araki et al. [56] and Fernandes et al. [43] by pulsed laser deposition.

Another important effect is the known loss of Sn in CZTS during the annealing process. Above 350°C Sn tends to re-evaporate and stoichiometry is lost. Different corrections of this effect have been presented among which the addition of inert gas in the chamber during the annealing or addition of SnS in the precursor to compensate the loss.

Weber et al. studied the loss of Sn in CZTS during the annealing process. It was found that desorption of SnS from the CZTS at temperature above 350◦C led to the losses of Sn in vacuum. The decomposition process of CZTS could be reduced by addition of an inert gas in the chamber where the sintering was carried out [44]. Redinger et al. suggested to add extra SnS and S materials in the precursor to prevent the decomposition of CZTS at high temperature [45].

Non vacuum deposition methods are very attractive since they would allow a lower pro‐ duction cost since expensive vacuum machines and high energy evaporation system would be avoided.

process where the single elements or binary chalcogenides compounds are vacuum-deposit‐ ed onto a heated substrate) obtaining a 2.3% efficiency cell. They also report, for the first time, that the electrical resistivity of the films can be improved considerably by treating the films with a KCN solution; this is still an important step of the fabrication process. Later on, Katagiri improved his deposition process by substituting Zn with ZnS in the e-beam se‐ quence and improved the adhesion of the absorber to the molybdenum, increasing its effi‐

An approach based on one-step synthesis of CZTS thin films by simultaneous coevaporation of precursor sources Cu, Zn, Sn, S was reported by Tanaka et al. The grain size of the film depended on the substrate temperature. Larger grains were obtained at higher substrate

An integrated vacuum apparatus, which combined the RF sputtering technique with the sul‐ furization process was employed to reduce the effect of moisture from air ambient on the property of CZTS films. Target sources were based on ZnS, SnS and Cu. Both the quality and reproducibility of the CZTS films were significantly improved by this method and a 5.7% conversion efficiency with the CZTS-based thin film solar cells was achieved [53]. By eliminating the metal oxide impurity in the CZTS films using deionized water, the efficiency

A lot of work has been recently presented by the IBM group; the research has been going on

For vacuum deposition a recent work from Wang et al. has been presented where the single elements are contemporarily evaporated on a SLG substrate by a rapid thermal deposition system. With an absorber thickness of only 1 micron and a few minutes annealing time the process is very promising for industrial application and has demonstrated a 6.8% conver‐ sion efficiency. High series resistance and a high charge recombination limit the efficiency [55]. Generally, if the composition of the precursor films is copper rich, it results in the forma‐ tion of a Cu2−x S secondary phase, which can be removed by a post treating with KCN

In metallic layer deposition it has been shown that the sequence of the different elements plays an important role in the crystallization of the absorber as shown with controversial re‐

Another important effect is the known loss of Sn in CZTS during the annealing process. Above 350°C Sn tends to re-evaporate and stoichiometry is lost. Different corrections of this effect have been presented among which the addition of inert gas in the chamber during the

Weber et al. studied the loss of Sn in CZTS during the annealing process. It was found that desorption of SnS from the CZTS at temperature above 350◦C led to the losses of Sn in vacuum. The decomposition process of CZTS could be reduced by addition of an inert gas in the chamber where the sintering was carried out [44]. Redinger et al. suggested to add extra SnS and S materials in the precursor to prevent the decomposition of CZTS at high temperature [45].

sults by Araki et al. [56] and Fernandes et al. [43] by pulsed laser deposition.

annealing or addition of SnS in the precursor to compensate the loss.

of the solar cells was further improved to 6.7% [49,54].

two different paths: vacuum and non vacuum deposition of CZTS.

ciency up to 1.08%.

100 Solar Cells - Research and Application Perspectives

temperature [52].

aqueous solution.

Like CdTe and CIGS also CZTS has been prepared by non-vacuum alternative methods such as spray pyrolysis, electrochemical deposition, and spin coating of precursor solutions, but in this case with very nice performance [48].

One of the first non vacuum synthesis of CZTS thin films by spray pyrolysis of precursor solutions was introduced in 1997 by Nakayama and Ito. The layer was fabricated by prepar‐ ing a mixture of water and ethanol with specific precursors such as CuCl, ZnCl2, SnCl4, and thiourea. With different precursors many other laboratories have studied this simple fabri‐ cation method, like Kumar et al. (high substrate temperature spray) [57] and Prabhakar et al. (ultrasonic spray) [58].

Spin-Coating of CZTS precursor solution is one of the most successful technique among the non vacuum methods but also among the vacuum ones. It is typically consisting in first the preparation of a precursor solution, then spin-coating the precursor solution on a SLG sub‐ strate and finally annealing the stacks in a controlled atmosphere to form the CZTS absorber.

Many different recipes have been presented from different researchers like Tanaka et al. [53], who presented an all-non vacuum deposited device. With a copper poor solution a 2.03% efficiency has been obtained. Other researcher like Pawar improved the crystallinity by the introduction of a complexing agent which increased inhomogeneity on the surface, later improved by sulfurization [59]. Fischereder et al. deposited CZTS films by spin coating a precursor solution consisting of metal salts of copper (I), tin (IV), and zinc (II) and thioace‐ tamide in pyridine. They found out that the formation of CZTS compound occurred at tem‐ perature as low as 105°C in vacuum. The band gap energy of the films varied between from 1.41 to 1.81 eV by changing the annealing temperature point [60]. This is probably due to the coexistence of secondary phases in the film under low annealing temperature.Todorov et al. firstly presented best efficiency by preparing a CZTS solution based on hydrazine. The CZTS films were obtained by annealing the precursor films in sulfur or sulfur and selenium atmosphere at high temperature. The best cell efficiency was 9.66%, later improved by Mitzi et al. up to 10.1%. These devices use a hybrid absorber which is selenized and sulfurized as well. Measurements presented and comparison with a typical high efficiency CIGS solar cell shows that this high performing device still suffers from dominant interface recombination, short minority carrier lifetime, and high series resistance.

In conclusions, CZTS is a new material, which has in the last ten years seen a huge improve‐ ment; a lot has been done to study the physical properties and to control the stoichiometry, especially secondary phases that are still a strong limitation to high efficiency. High series resistance and short minority carrier lifetime generally reduce the current of these devices and the tendency to form a great numbers of detrimental defects decreases the open circuit voltage.

However efficiencies above 10% have been demonstrated, moreover these devices can be successfully prepared in a very large variety of deposition methods from vacuum growth (coevaporation, sputtering, electron beam of precursors) to non-vacuum techniques among which the record cell has been obtained.

## **5. Conclusions**

The need for more and more energy supplies due to increased demand from emerging countries such as India, China and Brazil and the contemporary necessity to preserve the en‐ vironment, has increased interest in the development of new technologies that make use of solar energy.

In particular photovoltaic solar energy, the direct conversion of solar energy into electricity by means of semiconducting materials, has had very strong development in the last 20 years.

The most important parameter that characterizes a photovoltaic device is the ratio between its conversion efficiency and its cost. A value less than 0.5 \$/Wp is considered competitive with the electricity obtained from fossil energy sources (Grid-parity). Although large prog‐ ress in Cu(In,Ga)Se2 solar cell scientific knowledge and process management has been ac‐ complished, as evidenced by the high module and cell efficiencies fabricated by many groups, the range of deposition and device options that have been developed are not suffi‐ cient to ensure the achievement of this objective. In particular, for CIGS-based devices the question is: what is needed to be done to guarantee that Cu(In,Ga)Se2 solar cell technology reaches its potential for large-scale production?

We could first respond by indicating how critical the need is for opportunities to develop new production technologies, including enhanced equipment for the deposition of the lay‐ ers as well as production processes based on well-known engineering models. In addition a new second generation of production process-control tools will have to be worked out.

Improved equipment and very good diagnostic can be directly turned in to higher through‐ put, yield, and performance. This results in a better conversion efficiency-cost ratio.

In the ThiFiLab, considerable effort in the direction of simplifying the production process has been made and the achieved results suggest that we are closer to the development of a high-yield production line of low-cost highly efficient modules.

A second essential question could be: what might be the innovations that are needed for lay‐ ing the groundwork for a new generation of thin-film Cu(InGa)Se2-based solar cells?

A lot of R&D can be made by studying alternative substrates such as the ceramic tiles pro‐ posed by the ThiFiLab that directly allow the buildings integration of photovoltaics (BIPV).

Another step may be achieved by developing low-temperature manufacturing processes that allows the use of innovative substrates such as polymeric ribbons, which represents a basic passage towards the low-cost Roll-to-Roll (R2R) technology. Similarly, there may be significant cost and processing advantages to a cell structure that enables the use of a Cu(In‐ Ga)Se2 layer much less than 1 μm.

R&D on Cu(In,Ga)Se2 remains very attractive as well as promising through all of these chal‐ lenges, aimed at improving the knowledge of both the material and the device, in addition to developing new production technologies and revolutionary innovations.

The great stability over time shown by the high-efficiency devices together with the great variability of material stoichiometry and manufacturing processes provide a great expecta‐ tion that this material will give its contribution to energy demand at a competitive price while respecting the environment.

In this groove the research on kesterites is to be placed; these are relatively new materials for solar cell applications, and a lot of research remains to be done before the best components and processing methods are found. In particular, concerning Cu2ZnSnSe4 as PV material a bigeffort must be made in the knowledge related to the segregation of phases, which af‐ fectthe final functioning of the device. The structure of this cell, being similar to that normal‐ ly used for CIGS solar cells, has already been well defined, but the deposition methods of the absorber layer are still subject of study. Despite this, with this material as absorber, solar cells with conversion efficiency greater than 10% have been realized.

## **Author details**

(coevaporation, sputtering, electron beam of precursors) to non-vacuum techniques among

The need for more and more energy supplies due to increased demand from emerging countries such as India, China and Brazil and the contemporary necessity to preserve the en‐ vironment, has increased interest in the development of new technologies that make use of

In particular photovoltaic solar energy, the direct conversion of solar energy into electricity by means of semiconducting materials, has had very strong development in the last 20 years.

The most important parameter that characterizes a photovoltaic device is the ratio between its conversion efficiency and its cost. A value less than 0.5 \$/Wp is considered competitive with the electricity obtained from fossil energy sources (Grid-parity). Although large prog‐ ress in Cu(In,Ga)Se2 solar cell scientific knowledge and process management has been ac‐ complished, as evidenced by the high module and cell efficiencies fabricated by many groups, the range of deposition and device options that have been developed are not suffi‐ cient to ensure the achievement of this objective. In particular, for CIGS-based devices the question is: what is needed to be done to guarantee that Cu(In,Ga)Se2 solar cell technology

We could first respond by indicating how critical the need is for opportunities to develop new production technologies, including enhanced equipment for the deposition of the lay‐ ers as well as production processes based on well-known engineering models. In addition a new second generation of production process-control tools will have to be worked out.

Improved equipment and very good diagnostic can be directly turned in to higher through‐

In the ThiFiLab, considerable effort in the direction of simplifying the production process has been made and the achieved results suggest that we are closer to the development of a

A second essential question could be: what might be the innovations that are needed for lay‐

A lot of R&D can be made by studying alternative substrates such as the ceramic tiles pro‐ posed by the ThiFiLab that directly allow the buildings integration of photovoltaics (BIPV).

Another step may be achieved by developing low-temperature manufacturing processes that allows the use of innovative substrates such as polymeric ribbons, which represents a basic passage towards the low-cost Roll-to-Roll (R2R) technology. Similarly, there may be significant cost and processing advantages to a cell structure that enables the use of a Cu(In‐

put, yield, and performance. This results in a better conversion efficiency-cost ratio.

ing the groundwork for a new generation of thin-film Cu(InGa)Se2-based solar cells?

high-yield production line of low-cost highly efficient modules.

which the record cell has been obtained.

102 Solar Cells - Research and Application Perspectives

reaches its potential for large-scale production?

Ga)Se2 layer much less than 1 μm.

**5. Conclusions**

solar energy.

Alessio Bosio1\*, Daniele Menossi1 , Alessandro Romeo2 and Nicola Romeo1

\*Address all correspondence to: alessio.bosio@unipr.it

1 Physics Dept. University of Parma, Italy

2 Computer Science Dept.Universityof Verona, Italy

## **References**


[26] Noufi, R. (1994). *Photovoltaic Insidrr's Report*, *I-*, 31, 2-4.

[7] Romeo, N., Sberveglieri, G., Tarricone, L., & Paorici, C. (1977). *Appl. Phys. Lett.*, 30,

[9] Devaney, W., Michelsen, R., & Chen, W. (1985). Paper presented at 18th IEEE Photo‐

[10] Tuttle, I. R., Ward, J. S., Berens, T., Duda, A., Contreras, M. A., Ramanathan, K. R., Tennant, A. L., Keane, J., Cole, E. D., Emery, K., & Noufi, R. (1996). Proc. Materials

[11] Contreras, M. A., Egaas, B., Ramanathan, K., Hiltner, J., Swarzlander, A., Hasan, F., &

[12] Ramanathan, K., Contreras, M. A., Perkins, C. L., Asher, S., Hasoon, F. S., Keane, J., Young, D., Romero, M., Metzger, W., Noufi, R., Ward, J., & Duda, A. (2003). *Progr.*

[13] Hedstroem, J., Olsen, H., Bodegard, M., Kyler, A., Stolto, L., Hariskos, D., Ruckh, M., & Schock, H. W. (1993). *Proc. 23rd IEEE Photovoltaic Specialist Conf.*, 364-371.

[14] Chen, W., Stewar, J. M., & Stanbery, B. J. (1987). *Proc. 19th IEEE Photovoltaic Specialist*

[16] Gabor, A. M., Tuttle, J. R., Bode, M. H., Franz, A., Tennant, A. L., Contreras, M. A., Noufi, R., Jansen, D. G., & Heramnn, A. M. (1996). *Sol. Energy Mater. Sol. Cells*, 4,

[17] Tarrant, D., & Ermer, J. (1993). *Proc. 23rd IEEE Photovoltaic Specialist Conf.*, 372-375.

[19] Noufi, R., Matson, R. J., Powell, R. C., & Herrington, C. (1986). *Solar Cells*, 16, 479.

[18] Mickelsen, R. A., Chen, W. S., Hsiao, Y. R., & Lowe, V. (1984). *IEEE Trans. Electr. De‐*

[20] Birkmire, R. W., Dinetta, L. C., Lasswell, P. G., Meakin, J. D., & Phillips, J. E. (1986).

[22] Piekoszewski, J., Loferski, J. J., Bealieu, R., Beall, J., Roessler, B., & Shewchun, J.

[23] Romeo, N., Canevari, V., Sberveglieri, G., Bosio, A., & Zanotti, L. (1986). *Solar Cells*,

[24] Samaan, A. N., Abdul-Karim, N., Abdul-Hussein, N., Tomlinson, R. D., Hill, A. E., &

[8] Kazmerski, L., White, F., & Morgan, G. (1976). *Appl. Phys. Lett.*, 29, 268-269.

voltaic Specialists Conference. IEEE Publishing, NY, 173.

Noufi, R. (1999). *Progr. Photovolt: Res. Appl.*, 7, 311.

Research Society,. San Francisco.

*Photovolt: Res. Appl.*, 11, 225.

[15] Potter, R. (1986). *Sol. Cells*, 16, 521-527.

*Conf.*, 1445-1447.

*vices*, 311984, 542.

*Solar Cells*, 16, 419.

[21] Potter, R. R. (1986). *Solar cells*, 16, 521.

(1980). *14th IEEE Photovoltaic Specilist' Conf. 98.*

Armour, D. G. (1980). *Jpn. J. Appl. Phys.*, 19-3.

[25] Stolt, L. (1993). *Photovoltaic Insider's Report,*, *XII-*, 6, 1-2.

247-260.

16, 155.

108.

104 Solar Cells - Research and Application Perspectives

