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Effects of residual copper selenide on CuInGaSe2 solar cells
Solid-State Electronics 56 (2011) 175–178
Contents lists available at ScienceDirect
Solid-State Electronics journal homepage: www.elsevier.com/locate/sse
Effects of residual copper selenide on CuInGaSe2 solar cells Tung-Po Hsieh ⇑, Chia-Chih Chuang, Chung-Shin Wu, Jen-Chuan Chang, Jhe-Wei Guo, Wei-Chien Chen Green Energy and Environment Research Laboratories, Industrial Technology Research Institute, Hsinchu 310, Taiwan, ROC
a r t i c l e
i n f o
Article history: Received 1 July 2010 Received in revised form 15 November 2010 Accepted 22 November 2010 Available online 14 December 2010 The review of this paper was arranged by Prof. E. Calleja Keywords: CuInGaSe2 (CIGS) CuSe Thin-film Solar cells
a b s t r a c t Large-grain, copper-poor CuInGaSe2 (CIGS) films are favored in the fabrication of highly efficient solar cells. However, the degradation of cell performance caused by residual copper selenide (Cu2 xSe) remains a problem. This work studies the formation and behavior of excess CuxSe and further compares the cell performance of typical copper-poor with that of copper-rich solar cells. Since excess Cu2 xSe cannot be exhausted during the growth, it fully surrounds the polycrystalline CIGS grains. Excess Cu2 xSe in the CIGS film produces serious shunt paths and causes the pn junction to be of poor quality. A short circuit in copper-rich CIGS solar cells is attributable to the conductive Cu2 xSe. The best way to ensure high-efficiency of the cells is to exhaust Cu2 xSe during growth. Otherwise, a dense, chemically treated CIGS film is required to prevent the negative effects of excess Cu2 xSe. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Thin-film solar cells have become an important area of research owing to economic, energy-related, and environmental issues. Interest in thin-film solar cells is motivated by the potential of producing them at low cost. Of all thin-film solar cells, CuInGaSe2 (CIGS) solar cells have the greatest potential for replacing present technologies because of the low cost of their production and high conversion efficiency. Currently, 19.9%-efficient CIGS solar cells have been produced by co-evaporation [1]. The large-grain growth in polycrystalline CIGS is important to the improvement of cell efficiency. As is widely believed, copper selenide (CuSe) formed by copper-rich CIGS is an excellent medium for enhancing the growth of grains on CIGS absorbers. Previous studies have established that the vapor–liquid–solid mechanism can explain the grain growth of CIGS [2,3]. Additionally, copper-poor CIGS film is required to fabricate highly efficient solar cells. The elimination of residual Cu2 xSe is critical to CIGS solar cells. Currently, the most effective process is three-stage co-evaporation, which maintains a Cu-rich environment during grain growth and exhausts residual copper selenides to yield high-performance solar cells [1,4–7]. Another approach, which is a post-chemical process, for selectively removing surface Cu2 xSe has been reported to transform copper-rich CIGS films into copper-poor ones [8]. The degradation caused by residual Cu2 xSe is still a problem when the growth of CIGS absorbers is not precisely controlled. The problem is more serious in other approaches, such as sputtering or printing, during selenization. Although some ⇑ Corresponding author. E-mail address: [email protected] (T.-P. Hsieh). 0038-1101/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.sse.2010.11.019
papers systematically discuss the physics of residual Cu2 xSe on CIGS [2,3], the electrical properties of Cu-rich CIGS solar cells with residual Cu2 xSe are seldom discussed. In addition to investigating the formation and behavior of residual Cu2 xSe on CIGS films, this study further analyzes the performance of copper-poor and copper-rich solar cells. Notably, cell performance can be improved when the residual Cu2 xSe is removed. This study also provides direct evidence of the short circuit path in CIGS films due to the residual Cu2 xSe. 2. Experiment The growth method we used herein was a typical three-stage process [1]. The polycrystalline CIGS absorbers were co-deposited by Cu, In, Ga, and Se elements on 0.5-lm-thick molybdenum (Mo), which was sputtered on pieces of soda-limed glass. In the first stage, a (InGa)2Se3 compound was deposited at a substrate temperature of 400 °C. Then, the substrate temperature was increased to 550 °C for the deposition of Cu and Se. In the final stage, (InGa)2Se3 was deposited at 550 °C. During the final minute of deposition, only the In2Se3 compound was deposited in the absence of Ga. The total Ga/(Ga + In) ratio of CIGS films was maintained at 0.2. The typical absorber thickness was about 2.2 lm. Samples were designed to study the effects of residual Cu2 xSe on CIGS absorbers. The amount of excess Cu2 xSe was controlled by adjusting the ratio of Cu2 xSe to (InGa)2Se3. Some as-grown samples were adopted in material analysis. The surface morphology of the as-grown CIGS films was observed using a scanning electron microscope (SEM). The crystal phases and material quality were analyzed by X-ray diffraction. To determine the device
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characteristics, a 75-nm-thick cadmium sulfide (CdS) buffer layer was deposited on the CIGS absorber by chemical bath deposition and then a 100-nm-thick intrinsic ZnO layer was deposited by radio-frequency (RF) sputtering. The window layer (Al:ZnO) and patterned metal grid (Ni/Al) were also deposited by sputtering. Fig. 1 presents the full structure of the solar cells. Current–voltage curves of solar cells were plotted under standard air mass 1.5 global (AM 1.5G) illumination at 25 °C.
3. Results and discussion Fig. 2a and b shows the top-view and cross-section SEM images of our as-grown copper-poor CIGS film. Fig. 2a presents a smooth surface and compact CIGS grains. Small crystallites with the triangular shape on the surface reveal a Cu-poor surface [9]. A low void density is observed to yield high-efficiency solar cells. According to Fig. 2b, the grain size exceeds 2 lm with few lateral grain boundaries. The transportation loss of carriers should be suppressed. X-ray diffraction reveals a clear CIGS (1 1 2) diffraction at 2h = 26.95° with a half width at half maximum of 2h at around 0.18°, as shown in the top spectrum of Fig. 3. The narrow linewidth of X-ray diffraction (XRD) indicates the (1 1 2) direction of polycrystalline CIGS grows well, which can explain the mentioned large and compact grains shown in Fig. 2. By the Vegard’s law fitting, the ratio of Ga/(In + Ga) is about 0.2, as determined from the position of the diffraction peak. The strong diffraction can be attributed to the large and compact polycrystalline grains. The spectrum peaks at approximately 2h = 40.4° and 44.8° correspond to the (1 1 0) diffraction of Mo and the (2 2 0)/(2 0 4) diffraction of CIGS, respectively. The result shows that the preferred orientation of this sample is the (1 1 2) direction. In contrast to the (1 1 2) orientation, (2 2 0)/(2 0 4)-oriented CIGS films may contain a lower density of nonradiative recombination centers [7]. Control of the ratio of preferred orientation of high-efficiency solar cells by varying Se flux has been discussed [10]. The effects of preferred orientation of CIGS absorbers are obvious when the cell efficiency is over 15%. In this case, control of the preferred orientation was not considered. During grain growth, excess CuSe was formed on copper-rich CIGS film to promote crystalline growth. In the three-stage process, the copper content of the CIGS is easily adjusted by tuning the growth time between the CuSe and (InGa)2Se3 stages. Fig. 3 also presents several X-ray diffraction spectra that correspond to CIGS films with various copper contents. The estimated Cu/(In + Ga) ratio ranges from 0.98 to 1.3. As the ratio of CuSe to (InGa)2Se3 in-
Fig. 2. (a) Top-view and (b) cross-section SEM images of typical three-stage coevaporation CIGS film.
Fig. 3. X-ray diffraction of various copper content of CIGS films deposited on Mocoated soda-lime glass.
Fig. 1. The layer structure of the three-stage co-evaporation CIGS solar cells.
creases, the second diffraction spectrum did not change markedly, because the film remained copper-poor. As the ratio increased, Cu2 xSe diffraction signals appeared, indicating that the films had been transformed from copper-poor to copper-rich films, as presented in the third and fourth spectra. The XRD peaks of copper selenide are very close to the peaks of CuInSe2. The previous papers offer a hint that the shoulder on the CIGS (1 1 2) peak is associated with the copper selenide [11,12]. The excess Cu2 xSe
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cannot be fully exhausted during the growth. In the transition from copper-poor to copper-rich, the surface changed continuously as the Cu(In, Ga)3Se5 compound was reduced and Cu2 xSe began to appear. The most commonly observed Cu2 xSe phase in the Cu-rich CIGS is Cu2Se, but other copper selenides may have formed. Since the environment is Cu-rich, liquid-phase CuxSe acts as a flux agent that improves grain growth. Although each single polycrystalline grain of CIGS remains large, the diffraction peak intensity of CIGS (1 1 2) declines rapidly, suggesting that the structure of the CIGS absorber has changed. Actually, the X-ray diffraction presents the mixture of the crystalline CIGS and Cu2 xSe phase. The Cu-rich film in Fig. 4a and b has larger and deeper holes than the Cu-poor film. The copper selenide was not fully consumed when the grain growth in the third stage was incomplete. The surface of the Cu-rich film was relatively rough and loose due to the residual copper selenide. Its surface roughness (RMS108 nm) exceeded that of the Cu-poor film (RMS50–70 nm). When the buffer layers, CdS and ZnO, were deposited on the CIGS films, the coverage could not be uniform because of the presence of pinholes. The pinholes cause junction leakage due to the bad coverage of CdS and i-ZnO. Some defects are easily formed close to the pn junction. If this film is adopted to fabricate a solar cell, serious shunt paths and a poor pn junction will be the main problems that affect cell performance, because of the presence of unfavorable voids. In Cu-rich samples, residual copper selenide was left in the CIGS absorber. It is known that a potassium cyanide (KCN) was used to selectively remove the Cu2 xSe phase from the as-grown films.
Fig. 4. SEM images of Cu-rich CIGS film (a), (b) before and (c), (d)after KCN treatment. [(a), (c): plane-view; (b), (d):cross-section] (e) X-ray diffraction of Curich CIGS film before and after KCN treatment.
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Twenty minutes of KCN etching on the sample shown in Fig. 4a and b should remove all of the copper selenide even that located at the grain boundary. The CIGS surface with some pits and pinholes is observed, as presented in Fig. 4c and d. The etched pits on the top of CIGS are separately and randomly distributed. This result indicates that the CIGS surface is not fully covered by the thick Cu2 xSe grains. The distribution result is consistent with the results of a previous micro-Raman investigation [13]. Notably, not only the number of holes is increased, but also the original pinholes are enlarged in the boundaries. According to Fig. 4e, the copper selenide peaks of XRD disappeared after KCN etching, indicating that the pinholes originate from the copper selenide. This result gives a hint that excess Cu2 xSe is present on the grain boundaries. Interestingly, the cross-section image of the Cu-rich absorber in Fig. 4b shows some other, instead of polycrystalline CIGS, materials. According to aforementioned discussion, the Cu-rich CIGS is most likely to contain residual Cu2 xSe. The KCN etching can selectively remove all of the Cu2 xSe, and the polycrystalline structure and grain boundaries are then clearly observed, as shown in Fig. 4d. Therefore, the materials should be Cu2 xSe. When the CIGS grains are loose, the residual Cu2 xSe fills their vacancies or their grain boundaries. The distribution of Cu2 xSe in Fig. 4b reveals that excess Cu2 xSe is present not only on the surface but also throughout the CIGS film. The residual Cu2 xSe that surrounds the CIGS grains may further promote both lateral and vertical growth until the excess Cu2 xSe is exhausted. According to a previous paper, residual Cu2 xSe is commonly segregated on the surface [8]. A similar observation is also demonstrated on CuInSe2/GaAs material system [14]. The result is attributable to the compactness of the CIGS grains. The residual Cu2 xSe easily floats on the surface because the CIGS grains are compact and well connected. In this investigation, the residual Cu2 xSe phase segregates not only on the surface but also at the grain boundary, because the CIGS grains are incomplete and loose. The CIGS growth is able to exhaust the copper selenide both on the grain boundary and surface by increasing the deposition time of the third stage. Although residual Cu2 xSe can be selectively etched by KCN after growth, the etching forms some voids at the grain boundary. The best way to produce highly efficient CIGS solar cells is to exhaust the residual Cu2 xSe during the growth. If chemical etching is required to remove the residual Cu2 xSe, then a compact CIGS absorber must be used to segregate the residual Cu2 xSe at the surface. The SEM images show that the excess Cu2 xSe may connect the top surface of CIGS absorber to the Mo contact. Cu2 xSe is known to be a conductive material because of its high doping concentration. Fig. 5 presents the current–voltage characteristics
Fig. 5. Current–voltage curves for (a) Cu-poor (Cu/(In + Ga) = 0.9), (b) slightly Curich (Cu/(In + Ga) = 1.1) and (c) heavy Cu-rich (Cu/(In + Ga) = 1.3) CIGS solar cells.
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4. Conclusions
Fig. 6. Current–voltage curves for Cu-rich CIGS solar cells with and without KCN treatment.
of the copper-poor and copper-rich CIGS solar cells by the solid and dashed curves, respectively. Accordingly, the cell structure obtained using copper-rich CIGS films is always short-circuited because of the excess Cu2 xSe, as presented in Fig. 5. The copper-poor cell with an active area of 0.4 cm2 has an efficiency of 11%. The curve yields an open-circuit voltage (Voc) of 0.52 V, a short circuit current density (Jsc) of 32 mA/cm2, and a fill factor (FF) of 0.67. The result is consistent with the mentioned mechanism. Fig. 6 may further explain the effects of residual Cu2 xSe on electrical properties. It shows two J–V curves for the Cu-rich CIGS solar cells with and without KCN treatment. The CIGS cell without KCN treatment (dashed line) exhibits a short-circuited nature, which is similar to the copper-rich CIGS cells shown in Fig. 5. The conductive Cu2 xSe may become the shunt paths, especially for a high Cu content. As the residual Cu2 xSe is etched selectively by KCN, the shunt paths should be removed. The J–V curve (solid line) for the etched sample presents a normal pn junction property except for a reverse-bias leakage current. The leakage current can be attributed to the abovementioned pinholes. Although the Cu-rich CIGS cells were treated by KCN, the J–V curve cannot show acceptable Voc, Jsc and conversion efficiency as Cu-poor devices owing to a poor CIGS/ CdS junction.
In summary, all of the Cu-rich CIGS absorbers (Fig. 4a and b) have a rougher surface with more pinholes than that of the Cu-poor CIGS absorbers (Fig. 2). The rough surface of CIGS with many pinholes would cause poor coverage of the CdS buffer layer. Highperformance CIGS solar cells possess a pn junction between a p-type CIGS absorber and a n-type heterojunction CdS. However, copper selenide is a conductive material. Therefore, the residual copper selenide on CIGS destroys the pn junction behavior due to the conductive copper selenide and the Cu2 xSe/CdS interface. Therefore, excess Cu2 xSe is responsible for the poor performance of solar cells because it produces serious pinhole defects, poor quality of the pn junction, and the formation of a ‘Cu2 xSe’ shunt path. Excess Cu2 xSe may exist on the surface and the grain boundary. Although the properties of Cu-rich CIGS absorbers can be improved by KCN treatment, the J–V curve cannot show as acceptable Voc, Jsc, and conversion efficiency as can Cu-poor devices. If chemical treatment is required to remove the excess Cu2 xSe, then a dense CIGS film is needed to prevent the negative effects of the residual Cu2 xSe. Acknowledgements The authors would like to acknowledge the support from Ministry of Economics Affairs, Taiwan. References [1] Repins I, Contreras MA, Egaas B, DeHart C, Scharf J, Perkins CL, et al. Prog Photovoltaics Res Appl 2008;16:235. [2] Wada T, Kohara N, Negami T, Nishitani M. J Mater Res 1997;12:1465. [3] Klenk R, Walter T, Schock H-W, Cahen D. Adv Mater 1993;5(2):114. [4] Ishizuka S, Yamada A, Matsubara K, Fons P, Sakurai K, Niki S. Appl Phys Lett 2008;93:124105. [5] Ishizuka S, Shibata H, Yamada A, Fons P, Sakurai K, Matsubara K, et al. Appl Phys Lett 2007;91:041902. [6] Ramanathan K, Teeter G, Keane JC, Noufi R. Thin Solid Films 2005;480:499. [7] Contreras MA, Romero MJ, Noufi R. Thin Solid Films 2006;511:51. [8] Niki S, Fons PJ, Yamada A, Lacroix Y, Shibata H, Oyanagi H, et al. Appl Phys Lett 1999;74:1630. [9] Dimmler B, Dittrich H, Schock HW. In: Proceedings of the 20th IEEE photovoltaic specialists conference. New York; 1988. p. 1426. [10] Chaisitsak S, Yamada A, Konagai M. Jpn J Appl Phys 2002;41:507. [11] Miyake H, Ohtake H, Sugiyama K. J Cryst Growth 1995;146:233. [12] Dhanam M, Manoj PK, Prabhu RR. J Cryst Growth 2005;280:425. [13] Witte W, Kniese R, Powalla M. Thin Solid Films 2008;517:867. [14] Fons P, Niki S, Yamada A, Oyanagi H. J Appl Phys 1998;84:6926.
Contents lists available at ScienceDirect
Solid-State Electronics journal homepage: www.elsevier.com/locate/sse
Effects of residual copper selenide on CuInGaSe2 solar cells Tung-Po Hsieh ⇑, Chia-Chih Chuang, Chung-Shin Wu, Jen-Chuan Chang, Jhe-Wei Guo, Wei-Chien Chen Green Energy and Environment Research Laboratories, Industrial Technology Research Institute, Hsinchu 310, Taiwan, ROC
a r t i c l e
i n f o
Article history: Received 1 July 2010 Received in revised form 15 November 2010 Accepted 22 November 2010 Available online 14 December 2010 The review of this paper was arranged by Prof. E. Calleja Keywords: CuInGaSe2 (CIGS) CuSe Thin-film Solar cells
a b s t r a c t Large-grain, copper-poor CuInGaSe2 (CIGS) films are favored in the fabrication of highly efficient solar cells. However, the degradation of cell performance caused by residual copper selenide (Cu2 xSe) remains a problem. This work studies the formation and behavior of excess CuxSe and further compares the cell performance of typical copper-poor with that of copper-rich solar cells. Since excess Cu2 xSe cannot be exhausted during the growth, it fully surrounds the polycrystalline CIGS grains. Excess Cu2 xSe in the CIGS film produces serious shunt paths and causes the pn junction to be of poor quality. A short circuit in copper-rich CIGS solar cells is attributable to the conductive Cu2 xSe. The best way to ensure high-efficiency of the cells is to exhaust Cu2 xSe during growth. Otherwise, a dense, chemically treated CIGS film is required to prevent the negative effects of excess Cu2 xSe. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Thin-film solar cells have become an important area of research owing to economic, energy-related, and environmental issues. Interest in thin-film solar cells is motivated by the potential of producing them at low cost. Of all thin-film solar cells, CuInGaSe2 (CIGS) solar cells have the greatest potential for replacing present technologies because of the low cost of their production and high conversion efficiency. Currently, 19.9%-efficient CIGS solar cells have been produced by co-evaporation [1]. The large-grain growth in polycrystalline CIGS is important to the improvement of cell efficiency. As is widely believed, copper selenide (CuSe) formed by copper-rich CIGS is an excellent medium for enhancing the growth of grains on CIGS absorbers. Previous studies have established that the vapor–liquid–solid mechanism can explain the grain growth of CIGS [2,3]. Additionally, copper-poor CIGS film is required to fabricate highly efficient solar cells. The elimination of residual Cu2 xSe is critical to CIGS solar cells. Currently, the most effective process is three-stage co-evaporation, which maintains a Cu-rich environment during grain growth and exhausts residual copper selenides to yield high-performance solar cells [1,4–7]. Another approach, which is a post-chemical process, for selectively removing surface Cu2 xSe has been reported to transform copper-rich CIGS films into copper-poor ones [8]. The degradation caused by residual Cu2 xSe is still a problem when the growth of CIGS absorbers is not precisely controlled. The problem is more serious in other approaches, such as sputtering or printing, during selenization. Although some ⇑ Corresponding author. E-mail address: [email protected] (T.-P. Hsieh). 0038-1101/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.sse.2010.11.019
papers systematically discuss the physics of residual Cu2 xSe on CIGS [2,3], the electrical properties of Cu-rich CIGS solar cells with residual Cu2 xSe are seldom discussed. In addition to investigating the formation and behavior of residual Cu2 xSe on CIGS films, this study further analyzes the performance of copper-poor and copper-rich solar cells. Notably, cell performance can be improved when the residual Cu2 xSe is removed. This study also provides direct evidence of the short circuit path in CIGS films due to the residual Cu2 xSe. 2. Experiment The growth method we used herein was a typical three-stage process [1]. The polycrystalline CIGS absorbers were co-deposited by Cu, In, Ga, and Se elements on 0.5-lm-thick molybdenum (Mo), which was sputtered on pieces of soda-limed glass. In the first stage, a (InGa)2Se3 compound was deposited at a substrate temperature of 400 °C. Then, the substrate temperature was increased to 550 °C for the deposition of Cu and Se. In the final stage, (InGa)2Se3 was deposited at 550 °C. During the final minute of deposition, only the In2Se3 compound was deposited in the absence of Ga. The total Ga/(Ga + In) ratio of CIGS films was maintained at 0.2. The typical absorber thickness was about 2.2 lm. Samples were designed to study the effects of residual Cu2 xSe on CIGS absorbers. The amount of excess Cu2 xSe was controlled by adjusting the ratio of Cu2 xSe to (InGa)2Se3. Some as-grown samples were adopted in material analysis. The surface morphology of the as-grown CIGS films was observed using a scanning electron microscope (SEM). The crystal phases and material quality were analyzed by X-ray diffraction. To determine the device
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characteristics, a 75-nm-thick cadmium sulfide (CdS) buffer layer was deposited on the CIGS absorber by chemical bath deposition and then a 100-nm-thick intrinsic ZnO layer was deposited by radio-frequency (RF) sputtering. The window layer (Al:ZnO) and patterned metal grid (Ni/Al) were also deposited by sputtering. Fig. 1 presents the full structure of the solar cells. Current–voltage curves of solar cells were plotted under standard air mass 1.5 global (AM 1.5G) illumination at 25 °C.
3. Results and discussion Fig. 2a and b shows the top-view and cross-section SEM images of our as-grown copper-poor CIGS film. Fig. 2a presents a smooth surface and compact CIGS grains. Small crystallites with the triangular shape on the surface reveal a Cu-poor surface [9]. A low void density is observed to yield high-efficiency solar cells. According to Fig. 2b, the grain size exceeds 2 lm with few lateral grain boundaries. The transportation loss of carriers should be suppressed. X-ray diffraction reveals a clear CIGS (1 1 2) diffraction at 2h = 26.95° with a half width at half maximum of 2h at around 0.18°, as shown in the top spectrum of Fig. 3. The narrow linewidth of X-ray diffraction (XRD) indicates the (1 1 2) direction of polycrystalline CIGS grows well, which can explain the mentioned large and compact grains shown in Fig. 2. By the Vegard’s law fitting, the ratio of Ga/(In + Ga) is about 0.2, as determined from the position of the diffraction peak. The strong diffraction can be attributed to the large and compact polycrystalline grains. The spectrum peaks at approximately 2h = 40.4° and 44.8° correspond to the (1 1 0) diffraction of Mo and the (2 2 0)/(2 0 4) diffraction of CIGS, respectively. The result shows that the preferred orientation of this sample is the (1 1 2) direction. In contrast to the (1 1 2) orientation, (2 2 0)/(2 0 4)-oriented CIGS films may contain a lower density of nonradiative recombination centers [7]. Control of the ratio of preferred orientation of high-efficiency solar cells by varying Se flux has been discussed [10]. The effects of preferred orientation of CIGS absorbers are obvious when the cell efficiency is over 15%. In this case, control of the preferred orientation was not considered. During grain growth, excess CuSe was formed on copper-rich CIGS film to promote crystalline growth. In the three-stage process, the copper content of the CIGS is easily adjusted by tuning the growth time between the CuSe and (InGa)2Se3 stages. Fig. 3 also presents several X-ray diffraction spectra that correspond to CIGS films with various copper contents. The estimated Cu/(In + Ga) ratio ranges from 0.98 to 1.3. As the ratio of CuSe to (InGa)2Se3 in-
Fig. 2. (a) Top-view and (b) cross-section SEM images of typical three-stage coevaporation CIGS film.
Fig. 3. X-ray diffraction of various copper content of CIGS films deposited on Mocoated soda-lime glass.
Fig. 1. The layer structure of the three-stage co-evaporation CIGS solar cells.
creases, the second diffraction spectrum did not change markedly, because the film remained copper-poor. As the ratio increased, Cu2 xSe diffraction signals appeared, indicating that the films had been transformed from copper-poor to copper-rich films, as presented in the third and fourth spectra. The XRD peaks of copper selenide are very close to the peaks of CuInSe2. The previous papers offer a hint that the shoulder on the CIGS (1 1 2) peak is associated with the copper selenide [11,12]. The excess Cu2 xSe
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cannot be fully exhausted during the growth. In the transition from copper-poor to copper-rich, the surface changed continuously as the Cu(In, Ga)3Se5 compound was reduced and Cu2 xSe began to appear. The most commonly observed Cu2 xSe phase in the Cu-rich CIGS is Cu2Se, but other copper selenides may have formed. Since the environment is Cu-rich, liquid-phase CuxSe acts as a flux agent that improves grain growth. Although each single polycrystalline grain of CIGS remains large, the diffraction peak intensity of CIGS (1 1 2) declines rapidly, suggesting that the structure of the CIGS absorber has changed. Actually, the X-ray diffraction presents the mixture of the crystalline CIGS and Cu2 xSe phase. The Cu-rich film in Fig. 4a and b has larger and deeper holes than the Cu-poor film. The copper selenide was not fully consumed when the grain growth in the third stage was incomplete. The surface of the Cu-rich film was relatively rough and loose due to the residual copper selenide. Its surface roughness (RMS108 nm) exceeded that of the Cu-poor film (RMS50–70 nm). When the buffer layers, CdS and ZnO, were deposited on the CIGS films, the coverage could not be uniform because of the presence of pinholes. The pinholes cause junction leakage due to the bad coverage of CdS and i-ZnO. Some defects are easily formed close to the pn junction. If this film is adopted to fabricate a solar cell, serious shunt paths and a poor pn junction will be the main problems that affect cell performance, because of the presence of unfavorable voids. In Cu-rich samples, residual copper selenide was left in the CIGS absorber. It is known that a potassium cyanide (KCN) was used to selectively remove the Cu2 xSe phase from the as-grown films.
Fig. 4. SEM images of Cu-rich CIGS film (a), (b) before and (c), (d)after KCN treatment. [(a), (c): plane-view; (b), (d):cross-section] (e) X-ray diffraction of Curich CIGS film before and after KCN treatment.
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Twenty minutes of KCN etching on the sample shown in Fig. 4a and b should remove all of the copper selenide even that located at the grain boundary. The CIGS surface with some pits and pinholes is observed, as presented in Fig. 4c and d. The etched pits on the top of CIGS are separately and randomly distributed. This result indicates that the CIGS surface is not fully covered by the thick Cu2 xSe grains. The distribution result is consistent with the results of a previous micro-Raman investigation [13]. Notably, not only the number of holes is increased, but also the original pinholes are enlarged in the boundaries. According to Fig. 4e, the copper selenide peaks of XRD disappeared after KCN etching, indicating that the pinholes originate from the copper selenide. This result gives a hint that excess Cu2 xSe is present on the grain boundaries. Interestingly, the cross-section image of the Cu-rich absorber in Fig. 4b shows some other, instead of polycrystalline CIGS, materials. According to aforementioned discussion, the Cu-rich CIGS is most likely to contain residual Cu2 xSe. The KCN etching can selectively remove all of the Cu2 xSe, and the polycrystalline structure and grain boundaries are then clearly observed, as shown in Fig. 4d. Therefore, the materials should be Cu2 xSe. When the CIGS grains are loose, the residual Cu2 xSe fills their vacancies or their grain boundaries. The distribution of Cu2 xSe in Fig. 4b reveals that excess Cu2 xSe is present not only on the surface but also throughout the CIGS film. The residual Cu2 xSe that surrounds the CIGS grains may further promote both lateral and vertical growth until the excess Cu2 xSe is exhausted. According to a previous paper, residual Cu2 xSe is commonly segregated on the surface [8]. A similar observation is also demonstrated on CuInSe2/GaAs material system [14]. The result is attributable to the compactness of the CIGS grains. The residual Cu2 xSe easily floats on the surface because the CIGS grains are compact and well connected. In this investigation, the residual Cu2 xSe phase segregates not only on the surface but also at the grain boundary, because the CIGS grains are incomplete and loose. The CIGS growth is able to exhaust the copper selenide both on the grain boundary and surface by increasing the deposition time of the third stage. Although residual Cu2 xSe can be selectively etched by KCN after growth, the etching forms some voids at the grain boundary. The best way to produce highly efficient CIGS solar cells is to exhaust the residual Cu2 xSe during the growth. If chemical etching is required to remove the residual Cu2 xSe, then a compact CIGS absorber must be used to segregate the residual Cu2 xSe at the surface. The SEM images show that the excess Cu2 xSe may connect the top surface of CIGS absorber to the Mo contact. Cu2 xSe is known to be a conductive material because of its high doping concentration. Fig. 5 presents the current–voltage characteristics
Fig. 5. Current–voltage curves for (a) Cu-poor (Cu/(In + Ga) = 0.9), (b) slightly Curich (Cu/(In + Ga) = 1.1) and (c) heavy Cu-rich (Cu/(In + Ga) = 1.3) CIGS solar cells.
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4. Conclusions
Fig. 6. Current–voltage curves for Cu-rich CIGS solar cells with and without KCN treatment.
of the copper-poor and copper-rich CIGS solar cells by the solid and dashed curves, respectively. Accordingly, the cell structure obtained using copper-rich CIGS films is always short-circuited because of the excess Cu2 xSe, as presented in Fig. 5. The copper-poor cell with an active area of 0.4 cm2 has an efficiency of 11%. The curve yields an open-circuit voltage (Voc) of 0.52 V, a short circuit current density (Jsc) of 32 mA/cm2, and a fill factor (FF) of 0.67. The result is consistent with the mentioned mechanism. Fig. 6 may further explain the effects of residual Cu2 xSe on electrical properties. It shows two J–V curves for the Cu-rich CIGS solar cells with and without KCN treatment. The CIGS cell without KCN treatment (dashed line) exhibits a short-circuited nature, which is similar to the copper-rich CIGS cells shown in Fig. 5. The conductive Cu2 xSe may become the shunt paths, especially for a high Cu content. As the residual Cu2 xSe is etched selectively by KCN, the shunt paths should be removed. The J–V curve (solid line) for the etched sample presents a normal pn junction property except for a reverse-bias leakage current. The leakage current can be attributed to the abovementioned pinholes. Although the Cu-rich CIGS cells were treated by KCN, the J–V curve cannot show acceptable Voc, Jsc and conversion efficiency as Cu-poor devices owing to a poor CIGS/ CdS junction.
In summary, all of the Cu-rich CIGS absorbers (Fig. 4a and b) have a rougher surface with more pinholes than that of the Cu-poor CIGS absorbers (Fig. 2). The rough surface of CIGS with many pinholes would cause poor coverage of the CdS buffer layer. Highperformance CIGS solar cells possess a pn junction between a p-type CIGS absorber and a n-type heterojunction CdS. However, copper selenide is a conductive material. Therefore, the residual copper selenide on CIGS destroys the pn junction behavior due to the conductive copper selenide and the Cu2 xSe/CdS interface. Therefore, excess Cu2 xSe is responsible for the poor performance of solar cells because it produces serious pinhole defects, poor quality of the pn junction, and the formation of a ‘Cu2 xSe’ shunt path. Excess Cu2 xSe may exist on the surface and the grain boundary. Although the properties of Cu-rich CIGS absorbers can be improved by KCN treatment, the J–V curve cannot show as acceptable Voc, Jsc, and conversion efficiency as can Cu-poor devices. If chemical treatment is required to remove the excess Cu2 xSe, then a dense CIGS film is needed to prevent the negative effects of the residual Cu2 xSe. Acknowledgements The authors would like to acknowledge the support from Ministry of Economics Affairs, Taiwan. References [1] Repins I, Contreras MA, Egaas B, DeHart C, Scharf J, Perkins CL, et al. Prog Photovoltaics Res Appl 2008;16:235. [2] Wada T, Kohara N, Negami T, Nishitani M. J Mater Res 1997;12:1465. [3] Klenk R, Walter T, Schock H-W, Cahen D. Adv Mater 1993;5(2):114. [4] Ishizuka S, Yamada A, Matsubara K, Fons P, Sakurai K, Niki S. Appl Phys Lett 2008;93:124105. [5] Ishizuka S, Shibata H, Yamada A, Fons P, Sakurai K, Matsubara K, et al. Appl Phys Lett 2007;91:041902. [6] Ramanathan K, Teeter G, Keane JC, Noufi R. Thin Solid Films 2005;480:499. [7] Contreras MA, Romero MJ, Noufi R. Thin Solid Films 2006;511:51. [8] Niki S, Fons PJ, Yamada A, Lacroix Y, Shibata H, Oyanagi H, et al. Appl Phys Lett 1999;74:1630. [9] Dimmler B, Dittrich H, Schock HW. In: Proceedings of the 20th IEEE photovoltaic specialists conference. New York; 1988. p. 1426. [10] Chaisitsak S, Yamada A, Konagai M. Jpn J Appl Phys 2002;41:507. [11] Miyake H, Ohtake H, Sugiyama K. J Cryst Growth 1995;146:233. [12] Dhanam M, Manoj PK, Prabhu RR. J Cryst Growth 2005;280:425. [13] Witte W, Kniese R, Powalla M. Thin Solid Films 2008;517:867. [14] Fons P, Niki S, Yamada A, Oyanagi H. J Appl Phys 1998;84:6926.