Research Article | Open Access
Jiang Liu, Da-Ming Zhuang, Ming-Jie Cao, Chen-Yue Wang, Min Xie, Xiao-Long Li, "Preparation and Characterization of Cu(In,Ga)Se2 Thin Films by Selenization of Cu0.8Ga0.2 and In2Se3 Precursor Films", International Journal of Photoenergy, vol. 2012, Article ID 149210, 7 pages, 2012. https://doi.org/10.1155/2012/149210
Preparation and Characterization of Cu(In,Ga)Se2 Thin Films by Selenization of Cu0.8Ga0.2 and In2Se3 Precursor Films
Se-containing precursor films with two different compositions were prepared by magnetron sputtering from and targets, and then were selenized using Se vapor. The effects of precursor composition and selenization temperature on the film properties were investigated. The results show that Se phase plays a critical role in film growth and electrical properties of CIGS films. The Cu-rich films exhibit different surface morphology and better crystallinity, as compared to the Cu-poor films. All the CIGS films exhibit p-type conductivity. The resistivity of the Cu-rich films is about three orders of magnitude lower than that of the Cu-poor films, which is attributed to the presence of p-type highly conductive Se phase.
Cu()Se2 (CIGS) thin films have received considerable attention in recent years because of their application in solar cells, which have many advantages in high conversion efficiency [1, 2], the possibility of low-cost production [3, 4], and long-term stability [5, 6]. The world-record efficiency for small-area CIGS-based thin film solar cells has recently surpassed 20% . One of the special qualities of the CIGS material is its variable band-gap . The bandgap of chalcopyrite CIGS absorber could vary from 1.04 eV in CuInSe2 to 1.68 eV in CuGaSe2 by substituting indium for gallium. The addition of small amounts of Ga could not only raise the bandgap to more suitably match the AM1.5 solar spectrum, more importantly, but could also improve the electrical properties of CIGS films [9–11]. However, the CIGS films with higher gallium concentration show a significant loss of efficiency relative to lower gallium content films . The optimal [Ga]/[In + Ga] ratio of CIGS absorber for high efficiency CIGS solar cell is considered to be 0.2~0.3 from a number of experimental results [13–15].
A variety of processing techniques have been developed to prepare high quality CIGS thin films. The two most reported processes are the coevaporation from elemental sources and the selenization of metallic precursors. The coevaporation process consists of three stages, which all need precise control over the deposition rate of each element. The selenization of metallic precursors is a two-stage process, which involves the deposition of Cu-In-Ga metallic precursors in the first step followed by their selenization using H2Se gas or Se vapor . The presence of low-melting metal indium in Cu-In-Ga metallic precursors could always result in rough morphologies of the CIGS films and loss of indium from films during the ramp-up and selenization stages [17, 18]. Moreover, the CIGS films by selenization of metallic precursors have a poor adhesion to substrate, due to a threefold volume expansion caused by the incorporation of Se into the metallic precursors during selenization . Few works have been done using the selenium-containing precursors to prepare CIGS films. In this study, we introduce Cu0.8Ga0.2 and In2Se3 as precursor films and avoid the direct addition of elemental indium. The precursors were then treated by selenization using Se vapor. The effects of precursor compositions and selenization temperature on the film properties were investigated.
High purity In2Se3 ceramic target and CuGa alloy target with Ga content of 20 at% were used. The size of the two targets is 360 mm × 80 mm. In the first step the CuGa and In2Se3 films were deposited on Mo-coated and bare soda-lime glass substrate by middle frequency (MF) magnetron sputtering of CuGa and In2Se3 targets. By adjusting the sputtering time and current of the two targets, the precursor films with different compositions could be prepared. Generally, relative proportion of metallic elements in CIGS films lies on that of the precursor films. The copper content of the final film and during the deposition process has a strong influence on film growth and CIGS solar cell properties [19, 20]. For the purpose of comparison, two sets of as-sputtered films were prepared. One was Cu-rich composition ([Cu]/[In + Ga] > 1) and another was Cu-poor composition ([Cu]/[In + Ga] < 1). The substrate holder inside the sputtering equipment could move back and forth along the short transverse direction of targets to keep the homogeneity of the as-sputtered films. The substrate temperature and working pressure during sputtering was kept constant at 200°C and 0.5 Pa, respectively. In the second step, the as-sputtered films were selenized with Se vapor in a three-chamber furnace which can be easily scaled up to provide a large-area in-line sequential CIGS process. Prior to heating, the chambers were evacuated and flushed two times with high purity nitrogen. The selenization time was held constant at 30 min and the selenization temperature varied from 400°C to 600°C.
Crystalline structures of the samples were examined by X-ray diffraction carried out with D/max-RB diffractometer (Rigaku) using Cu Ka radiation (0.1541 nm). The compositions of the films were tested by X-ray fluorescence (LAB CENTER XRF-1800). The surface and cross-section micrographs of the films were obtained by thermal field emission scanning electron microscope (TFE-SEM, LEO-1530). Raman spectra were recorded using Renishaw Raman microscope (model RM2000) with an argon ion laser wavelength of 514.5 nm. Hall-effect measurements were used on films grown on bare soda-lime glass to determine the resistivity and carrier concentration of CIGS films.
3. Results and Discussion
The XRD patterns of the precursor films deposited on bare soda-lime glass with two different compositions are presented in Figure 1. The corresponding composition data are shown in Table 1. To investigate the relative compositional change, the corresponding atomic ratios of Cu/(In + Ga), Ga/(In + Ga), and 2Se/(Cu+3(In + Ga)) were also calculated. The films were intentionally deposited into two groups, Cu-rich and Cu-poor precursor films, which have [Cu]/[In + Ga] of 1.277 and 0.935, respectively. Since Ga is incorporated in CuGa alloy target, the Cu-rich precursor film deserves to have a slightly higher [Ga]/[In + Ga] than the Cu-poor precursor film. The 2[Se]/([Cu]+3[In + Ga]) of the two precursor films are far less than that of near-stoichiometric Cu()Se2 compound and need be improved during selenization. Figure 1 indicates that XRD patterns of the two precursor films are almost the same. The diffraction peaks at = 26.6° corresponds to (112) plane of CIGS chalcopyrite structure. Two broad peaks at around = 44° are attributed to Cu (111) and CIGS (204/220). No In-Se related peaks or other characteristic peaks are found, which could be explained by the low crystallinity of the films. After selenization at 550°C for 30 min, XRF measurements (Table 1) show that the CIGS films selenized from Cu-rich and Cu-poor precursor films still conserve their overall Cu-rich and Cu-poor compositions, respectively, with corresponding [Cu]/[In + Ga] of 1.293 and 0.938. The [Ga]/[In + Ga] ratios of the CIGS films are nearly equal to these of their respective precursor films, which could draw the conclusion that the element concentrations of indium and gallium kept almost constant during selenization. Additionally, the 2[Se]/([Cu]+3[In + Ga]) ratios improved noticeably after selenization, which of Cu-rich film is higher than that of Cu-poor film due to the monovalence chemical state of copper.
Figure 2 shows the XRD patterns of CIGS films selenized from Cu-rich precursor films at different selenization temperatures for 30 min. All samples show characteristic peaks of the chalcopyrite-type structure and are oriented along the (112) direction parallel to the substrate. The full width at half maximum (FWHM) of (112) peaks of the films decreases with the increase of selenization temperature, which may be attributed to the crystalline quality improvement and grain growth. For the CIGS film selenized at 550°C, two additional peaks at = 31.4° and 55.9° could be found, which correspond to MoSe2 (100) and MoSe2 (110), respectively, according to Joint Committee on Powder Diffraction Standards (JCPDS) Number 29-0914. The intensity of the two peaks of MoSe2 gets stronger when the selenization temperature increases from 550°C to 600°C, which is indicative of higher degree of selenization. Note that when the selenization temperature was 600°C, the soda-lime glass substrate warped but the film deposited on it still did not peel off and could be used for XRD analysis and other tests.
Figure 2(b) shows the detailed XRD patterns of CIGS films from to 47° where (204/220) peak of chalcopyrite structure lies. Overall, the position of the (204/220) peak seems to shift to higher diffraction angle as the selenization temperature increases. The Cu()Se2 solid solution is based on CuInSe2. Gallium is added to obtain CIGS with a wider bandgap; accordingly, the lattice constants, a and c, decrease. With the increasing of [Ga]/[In + Ga] ratio, the diffraction peaks of chalcopyrite phase shift to higher side. Theoretically, if the lattice constant c/a ratio does not equal two, the (204) and (220) peaks would uncouple. For the CIGS film selenized at 400°C, the splitting of the (204) and (220) peaks is observed, which is in very good agreement with the CuInSe2 phase (JCPDS Number 40-1487). For the CIGS film selenized at 550°C or 600°C, the (204) and (220) peaks almost overlap, the positions of which match well with the CuIn0.7Ga0.3Se2 solid solution (JCPDS Number 35-1102). Based on the above experimental observations, it could be concluded that CuInSe2 may form first at a relatively low selenization temperature of 400°C, or the formation of CuInSe2 occurs faster than the formation of CuGaSe2, which is also observed in other literatures [21–23]. It could be easily understood that the selenization temperature has a strong influence on the reaction rate and atomic diffusion coefficient. With the increasing of selenization temperature, the gallium diffuses into CuInSe2, thus resulting in the formation of CuIn0.7Ga0.3Se2 solid solution.
To further investigate the structural properties of CIGS thin films, Raman spectroscopy was adopted. Raman spectra of CIGS films selenized at different selenization temperatures for 30 min from Cu-rich precursor films and Cu-poor precursor films are presented in Figure 3. All spectra are normalized to the peak at around 172–174 cm−1, which corresponds to the dominant A1 mode of CIGS films resulting from the vibration of the Se anions in the plane with the cations at rest [24, 25]. And mixed B2/E modes are also observed at around 213 cm−1, which represent vibrations of anions and cations together . For all the Cu-rich films (Figure 3(a)), an additional mode appears at around 258–260 cm−1, which is assigned to the A1 mode of Cu-Se compounds like CuSe or Cu2Se and is labeled with in the following literatures [19, 26]. For the Cu-poor films (Figure 3(b)), the peak of A1 mode exists only when the selenization temperature was 400°C. From the above observation, it is clear that inadequately reaction for the Cu-poor films at low selenization temperature could result in phase segregation. For the Cu-rich films, even at high selenization of up to 600°C, segregation is always present, which could be explained with the Cu2Se-In2Se3 pseudo-binary phase diagram , where excess Cu couldn’t be incorporated in the chalcopyrite phase.
Surface SEM images of CIGS films selenized from Cu-rich precursor films and Cu-poor precursor films at different selenization temperatures are presented in Figure 4. The insets show the corresponding cross-section micrographs of the films. From Figure 4(a–e) and Figure 4(f–j), the crystallinity of the films improves noticeably as the selenization temperature increases. The Cu-rich films exhibit a granular surface morphology, whereas a triangular surface morphology is observed obviously for the Cu-poor films selenized at 550°C or 600°C. Low selenization temperature of 400°C results in a significant increase in surface roughness of the Cu-rich and Cu-poor films. In comparison with the Cu-poor films, the Cu-rich films have a larger grain size. When the selenization temperature was 500°C or above, the surface morphology of Cu-rich films exhibit closely packed grains which are approximate 1 μm wide. From the cross-section micrographs, the comparison of film growth between the Cu-rich films and the Cu-poor films is also evident. According to the growth model proposed by Klenk et al. , a low-temperature liquid Cu-Se phase is present for the Cu-rich films during selenization and can aid the film growth. Consequently, the difference of Cu content between two films could result in different morphology and grain size. Furthermore, the grain is layered along the cross-section direction, which is more obvious at relatively low selenization temperature. The grain size of the top layer is larger than that of the bottom layer. It could be explained with the degree of selenization. Higher selenization temperature could result in higher degree of selenization , which is also consistent with the XRD analysis. Since the reaction proceeds from the surface towards the bulk during selenization, the crystalline quality of the bottom layer increases as the selenization degree increases.
Figure 5 shows the resistivity and carrier concentration of CIGS films selenized from Cu-rich precursor films and Cu-poor precursor films as a function of selenization temperature. All the CIGS films exhibit p-type semiconductor material. In comparison with the Cu-poor films, a significant decrease in resistivity and increase in carrier concentration by several orders of magnitude are observed in the Cu-rich films, which can be explained by the electrical properties of p-type phase. The is a p-type highly conductive phase, could produce a large number of holes, and thus decreases the resistivity of bulk film. It should be noted that Cu-rich films with such high carrier concentration, which is beyond the optimum range of 1016-1017 cm−3 , cannot be used for devices without further treatment. In both cases of the Cu-rich and Cu-poor films, the resistivity and carrier concentration tend to increase and decrease, respectively, as the selenization temperature increases. Note that the electrical properties of CIGS films are dominated by native defects resulting from deviations from stoichiometric composition, such as vacancies and interstitials . Considering the SEM images presented in Figure 4, increasing selenization temperature results in the improvement of crystallinity and therefore in a decreased total number of defects, which may be related to net carrier concentrations of CIGS films.
A two-stage method has been developed to prepare Cu(In,Ga)Se2 films using Se-containing precursors in Se-vapor selenization process. The effects of precursor composition and selenization temperature on the film properties were investigated. The [Cu]/[In + Ga] and [Ga]/[In + Ga] ratios in CIGS films could be controlled by adjusting the sputtering parameters of the Cu0.8Ga0.2 and In2Se3 targets. Two sets of films, Cu-poor and Cu-rich films, were intentionally prepared for comparison. It is revealed that the amount of Cu in the precursor films strongly influences the morphology, structural and electrical properties of the subsequent selenized films. XRD analysis shows that Cu(In,Ga)Se2 solid solution needs higher selenization temperature due to the inclusion of gallium. For Cu-rich films, segregation is always present, but for Cu-rich films, it exists only when the selenization temperature was 400°C due to inadequate reaction. In comparison with the Cu-poor films, the Cu-rich films have a higher crystallinity due to the presence of low-temperature liquid Cu-Se phase. The crystallinity of the films improves noticeably as the selenization temperature increases. The resistivity of the Cu-rich films is about three orders of magnitude lower than that of the Cu-poor films, which is attributed to the presence of p-type highly conductive phase. In both cases of the Cu-rich and Cu-poor films, the dependence of resistivity on selenization temperature is similar and may be due to the change in film crystallinity and native defects.
The authors are thankful to Song Jun for equipment maintenance, Li Jun for XRF measurement, and Chen Feng-en for measurement of Raman spectra. This work was financially supported by Initiative Scientific Research Project of Tsinghua University.
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