The effect of Na in polycrystalline and epitaxial single-crystal CuIn 1 x Ga x Se 2

Transcription

1 Thin Solid Films (2005) The effect of Na in polycrystalline and epitaxial single-crystal CuIn 1 x Ga x Se 2 A. Rockett* University of Illinois, ESB, MC W. Springfield Ave., Urbana, IL 61801, United States Available online 15 December 2004 Abstract Na is found to improve the performance of Cu(In,Ga)Se 2 (CIGS) solar cells although the mechanism is not clear. This paper briefly reviews some of the observations on Na in CIGS polycrystalline and epitaxial films. Experiments suggest weak electrical effects of Na within grains, primarily by reducing compensation and in some cases by enhancing acceptor concentrations. As it segregates to surfaces, it has been suggested that Na acts through passivation of grain-boundary defects. However, its main effect is on device open-circuit voltage (and somewhat on fill factor), which does not correlate with grain size but rather with bulk grain defects. The Na concentration scales somewhat with grain boundary density averaged over large areas of film, suggesting that it may be active there. Modest Na concentrations often increase grain size in polycrystals, although not when the grain size is already large, and often changes preferred orientation. Na segregates to the surfaces of CIGS grains. These results suggest that it may act at the surface, modifying growth mechanisms or defect organization during growth. TEM evidence shows that strong concentration of Na in the grain boundaries, sufficient to passivate surface defects by itself, is unlikely to occur. Finally, Na is removed from the surface of CIGS by aqueous solutions, such as those used to form the heterojunction. It is concluded that Na acts at the surface during growth to organize point defects, probably including by reduction of vacancy populations, within the bulk grains but that it has no residual effect once growth is complete. D 2004 Elsevier B.V. All rights reserved. Keywords: Open-circuit voltage; Cu(In,Ga)Se 2 ; Na concentration 1. Introduction Recent environmental and energy resource concerns have increased interest in renewable energy sources, such as photovoltaic devices. CuIn 1 x Ga x Se 2 (CIGS)/CdS heterojunction devices are promising candidates having the highest efficiency, exceeding 19%, of any thin film polycrystalline solar cell [1]. The diodes work well when fabricated from polycrystalline materials and no evidence of a correlation between grain size and device performance is observed until the grains are much less than 1 Am in diameter [2]. Typically, one finds that the efficiency of the devices is improved significantly by the presence of Na during growth [3], as discussed below. The mechanism for this improvement is not agreed to this point. To complicate the picture, many * Tel.: ; fax: address: arockett@uiuc.edu. experimental results on the effect of Na appear either directly contradictory or to disagree in broad terms. This paper argues that the literature is not in fact contradictory but rather that the results differ because of the details of the experiments performed. Based on the existing literature and recent results, it appears that Na must act during film growth to improve crystal quality, rather than having a specific ongoing effect in devices. Consistent with this, there is no clear evidence of a persistent effect of Na itself on bulk grains, grain boundaries, or the heterojunction in the film, nor is there an obvious site at which Na acts within the films. 2. Electronic effects of NA The effect of Na on devices has been studied by numerous groups, primarily on materials produced by multisource evaporation [3 10]. Typical results show as /$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi: /j.tsf

2 A. Rockett / Thin Solid Films (2005) much as a 50% increase in device efficiency with Na addition, decreasing at both higher and lower concentrations (see, for example, Refs. [8,10]). The primary improvements are in the open-circuit voltage (V OC ) and often in fill factor (FF) with little or no change in collected current. The lack of a grain-size effect and analysis of both single-crystal and polycrystalline devices suggests that V OC is largely determined by recombination in the space-charge region; thus, an increase in V OC indicates a reduction in recombination centers or an increase in carrier concentration or both within the grains near the surface of the film with the addition of Na. In typical devices grown on soda-lime glass substrates coated with Mo, the Na is incorporated into the Mo back contact in the form of Na 2 O from the glass and diffuses into the device [11]. However, there is a strong preference for Na to segregate to the surface of CIGS films, indicating that Na incorporation into CIGS is energetically unfavorable [12 15]. Although it leaves the glass as Na 2 O, it is probably converted at least mostly to a selenide compound in the CIGS layer or on the surface due to the high activity of Se and the low activity of O during deposition process. Several groups have studied the electronic effect of Na within CIGS grains by Hall effect [16,17] and photoluminescence [17,18]. The results show no correlation of an effect of Na with the presence of grain boundaries nor do they show a difference in the effect of Na within for singlecrystal as opposed to polycrystalline materials [17]. Although Na has been suspected of doping CIGS, none of these studies showed any evidence of new states in the energy gap. Schroeder and Rockett [16] and Kimura et al. [18] found a reduction in compensation in CuInSe 2 (CIS) epitaxial and polycrystalline thin films, respectively, with no increase in carrier concentration while Schuler et al. found both reduction in compensation and increase in carrier concentration associated with Na addition in both polycrystalline and epitaxial CuGaSe 2 (CGS). In addition, most device results have been interpreted as having improved voltages resulting strictly from an increase in hole concentration in the absorber. While the effect on the majority carrier states varies, it seems generally agreed that Na reduces compensating defects and usually enhances carrier concentrations. The absence of an identifiable new state associated with Na also indicates that it is not a dopant in CIGS of any sort. In addition to direct measurements of the effect of Na on electrical properties of films, an additional indirect indication may be found in comparison of transient photocapacitance spectroscopy [19,20] and cathodoluminescence [21] depth profiling for epitaxial films without Na and polycrystals with Na, both of which showed identical deep defects and band tails in films with and without Na. Reduction in donor compensation without creation of a new acceptor state shows that rather than adding new point defects, Na is removing or reducing the donor defect responsible for compensation. This is consistent with the observed increase in carrier concentration in many devices. Reduction in compensation must be due to direct insertion into the defect causing passivation (in the case of a vacancy) or elimination of the defect (for either a vacancy or an antisite). The two most likely sources of the compensating donor are Se vacancies (V Se ) and In on Cu antisites (In Cu ). Inserting a group I alkali metal on a group VI lattice site seems unlikely and it is not obvious how the Na atom would passivate a Se vacancy directly. Rather, it seems probable for either donor that the Na acts to remove the defect rather than somehow deactivating it. Na has been found by several researchers [3,4] to reduce the resistivity of polycrystalline CIGS films. This could be due to an increase in acceptor density as observed in some films, or could be due to a reduction in grain boundary resistivity. The more likely explanation in this case is the latter as in-plane resistivity is probably dominated by grain boundary terms [17]. Reduced boundary resistivity could be due either to the presence of a conductive phase in the boundaries or to reduction in carrier depletion in the bulk grains around the boundaries. No conductive phase has been observed to be connected with the presence of Na nor has any phase in the boundaries been directly observed by highresolution transmission electron microscopy (TEM) [22,23]. A high-resolution TEM image of a typical grain boundary in a material produced at the University of Delaware Institute for Energy Conversion as described in Ref. [2] is given in Fig. 1 [23], showing a very abrupt interface with no evidence of any interfacial phases or distortion of lattice planes near the boundary. Fig. 1. A high-resolution transmission electron micrograph [21] showing the structure of a grain boundary in a multisource-evaporated polycrystalline CIGS film [9]. This is a high-angle grain boundary so lattice image contrast is not obtained in both grains. However, fine fringes are visible in the right-hand grain from which it can be seen that no distortion exists up the boundary. The boundary is sharp (roughly one monolayer wide) and there is no evidence of a chemical or structural change present.

3 4 A. Rockett / Thin Solid Films (2005) 2 7 Fig. 2. This figure shows the composition of grains (filled points) and grain boundaries (open points) for two samples obtained from Shell Solar (triangles) and Global Solar Energy (squares) for baseline device materials (after Ref. [20]). It seems most likely given the structure of the boundaries that the change in conductivity is the result of a change in boundary defects pinning the Fermi energy, or bulk defects within the grains. Certainly, an increase in conductivity in the grains might reduce the effective resistance of grain boundaries in the same way that doping can effectively convert a Schottky barrier to an ohmic contact due to tunneling breakdown of the barrier. One would expect that a direct and measurable effect within the grain boundaries might be correlated with a measurable deviation of the grain boundary composition relative to that of the bulk grains, considering the number of dangling bonds that boundaries normally include. Such a deviation has been inferred but not observed to date. Indeed, recent microchemical and microstructural characterizations show the contrary no difference in grain and grain boundary composition [22,23]. A typical comparison of the grain and grain boundary compositions for CIGS polycrystalline device layers measured by nanoprobe EDS is shown in Fig. 2. (The intrinsic scatter in the data in this figure is roughly F3 at.%, much less than the total scatter. A large amount of similar but unpublished data also exists, indicating a similar lack of difference in grain boundary composition [23]). This result suggests that any change in conductivity due to the presence of Na is the result of redistribution of atoms to reduce defects in the boundary and/or within the grains rather than because of a large change in boundary chemistry. 3. Structural effects of NA Several investigators have shown an increase in grain size in CIGS polycrystals grown by multisource evaporation as a result of addition of Na prior to growth [5,6,10,18,24], although another study [25] found a decrease in grain size with added Na. In addition, Bodeg3rd et al.[5,10] found an increase in (112) preferred orientation in the crystallites which has subsequently been observed by others. Although the results differ, the experiments also changed, especially in the form of delivery of the Na. Furthermore, the grains in Rudmann et al. [25] were large even in the absence of added Na. Thus, although not a ubiquitous observation, the increase in grain size and change in preferred orientation appears to be consistent throughout the preponderance of the experiments reported to date while differing results are probably valid but were produced in noncomparable experiments. Interpretation of the effects of Na on microstructure are further complicated in multistep processes where massive changes in stoichiometry and recrystallization of the film in general are occurring during growth. Nonetheless, a strong implication of an increase in grain size with the addition of Na for an otherwise constant process is either that the nucleation process is being altered or that the diffusivity of atomic species is being enhanced, especially on grain surfaces where growth occurs. A change of nucleation could account for the differing results and is also consistent with the change in grain orientation in many studies. A close-packed-plane-up orientation [(112) for chalcopyrite] is what one would expect for growth of any material on a surface having no preference for orientation the close-packed plane brings adatoms closest together and is widely observed in thin films of all sorts grown on amorphous substrates. Thus, a decreased interaction with the substrate could enhance such preferred orientation. Alternately, a study of the surface morphology and growth of epitaxial CIGS films has strongly suggested that the (112) surfaces have the lowest energy [26,27]. Thus, enhanced atomic diffusion would also be expected to favor this surface. The observations available to date are therefore not sufficient to establish the change in growth mechanism resulting from Na addition. However, it appears to be a surface-related phenomenon. Furthermore, Na does not appear to enhance diffusivities in the bulk material so the changed nucleation model seems more probable. The argument that Na changes growth modes by action upon adatoms on the surfaces of grains rather than in the bulk of the grains is supported by several studies showing a strong tendency of Na to segregate to the surface of bulk single crystals [15], epitaxial layers [16], and polycrystalline films [14]. Likewise, there is some indication that the concentration of Na is coupled to grain boundary density [8,28], suggesting that segregation is not limited to surfaces but may include rejection from grains to other interfaces, such as boundaries [29]. It has been proposed that Na, being isovalent with Cu, could form a chalcopyrite surface phase or related compound [30]. Although this phase is not thermodynamically stable as a bulk material, it could exist in thin layers on chalcopyrite surfaces. Evidence for such a compound has

4 A. Rockett / Thin Solid Films (2005) been found under certain conditions [31]. However, in typical completed devices finished with a dip-coated CdS layer forming the heterojunction, there is no evidence from secondary ion mass spectrometry (SIMS) analysis of a significantly high Na concentration at the heterojunction or any correlation of residual interfacial Na with device performance. For example, Fig. 3 shows a typical profile for Na in a CIGS/CdS heterojunction. The CIGS was deposited by a multilayer multisource evaporation process. The profile is not surprising as experimental evidence shows that Na is effectively removed from the surface of CIGS polycrystals by a water rinse [32]. The aqueous environment of the dip process would therefore be expected to remove the majority of any surface Na prior to finishing the heterojunction. Therefore, although surface segregation is widely observed, the effect of the Na does not appear to be associated with its presence in the junction. We may then turn to the grain boundaries for evidence of an effect of Na. One of the longstanding questions with regard to CIGS is how the devices can operate satisfactorily in the presence of grain boundaries to the point of showing very little change in efficiency as a function of grain size until the grains are very small [2,9]. The failure of epitaxial layers and single crystals to produce good devices also supports the idea that grain boundaries actually help the device performance. The question is, how? As noted above, Na concentration is often found to be correlated with grain size, although clear evidence of Na in the grain boundaries themselves is limited [29]. It has been proposed [33] that Na acts to passivate grain boundaries, possibly by bringing O with it and that this eliminates Se vacancies in the Fig. 3. A typical SIMS profile for a CdS/CIGS heterojunction showing the absence of significant accumulation of Na at the heterojunction. Indeed, there appears to be no more Na present at the junction than elsewhere in the film. For this chalcogen/chalcogen interface, no significant ion yield changes for Na are expected. boundaries. In a recent series of tests, Li et al. [22] and Lei et al. [23] examined individual grains and grain boundaries in CIGS from complete polycrystalline device layers produced by a variety of techniques at both commercial and academic laboratories. The studies involved microenergy dispersive X-ray analysis (A-EDS) carried out in a high-resolution analytical scanning transmission electron microscope (STEM). The probed region diameter was ~1 nm. The results showed no change in grain boundary composition relative to the bulk grains (Fig. 2) and no correlation of Na or O signal with grain composition. Na concentration was linearly related to O concentration, although the stoichiometry suggested by the slope of the relationship would be roughly NaO 2, indicating that some of the O may have been bonding to the matrix. I note that although this appears to be an odd stoichiometry for a Na compound, it is similar to that of Na selenides observed by Braunger et al. [28]. These authors also observed a correlation between Na and O concentration by SIMS in air-exposed films. This result supports the proposition that Na is simply bringing O into the structure and that O is replacing missing Se. However, this is ruled out by the fact that device performance is not correlated with O concentration in the grain boundaries [8], that the grain and grain boundary compositions are identical to within the error in the measurements [22], and that the measured stoichiometry of the grains and grain boundaries does not suggest a S+Se deficiency (Fig. 2). In short, there is no evidence from the STEM analysis of a connection of Na to the grain boundaries or to compensation for any nonstoichiometry either by itself or through added O. To summarize the arguments to this point: Na does not act in the heterojunction to improve performance as no significant excess of Na is present there and it would be expected to be removed during junction formation in any case. Na does not act in large quantities at the grain boundaries either directly or through incorporation of O as what little is present of either element is not correlated with stoichiometry there is no detectable difference in stoichiometry between the bulk grains and grain boundaries that would suggest a need for another element to be present and no separate phase is observable in the boundaries. Finally, Na segregates strongly out of the bulk of grains and does not dope the grains when it is present, suggesting that it is not acting on its own within the grains to improve the devices. Then the question arises, where is the Na that is certainly observed in the films? Probably some of it is in the grains and some in the grain boundaries. Certainly, too little is present in the grain boundaries to detect in EDS measurements (about 5% of the boundary sites); therefore, it is not correcting a major stoichiometry problem itself. Even less is present in the grains, less than would be needed to account directly for the change in the density of holes by substitional doping with Na (nor is it found to be a dopant). This leaves only one possible explanation for the effect of Na that is consistent with all of the experimental data that Na acts

5 6 A. Rockett / Thin Solid Films (2005) 2 7 during growth of the material upon the CIGS matrix material itself (as also proposed in Ref. [29]), and that after growth is complete, Na has no further direct effect. To test this, my group has begun to study the effects of Na on the structure growth mode of epitaxial CIGS layers and to determine the resulting change in device performance [34]. Details will be presented elsewhere but some of the initial results are summarized here. We observe, as have others, that Na reduces the diffusivity of Ga in CIGS. In my group s results, the effect is reflected in apparently reduced diffusion of Ga from the GaAs substrate into the CIS epitaxial layers and is quite pronounced. In our results, the reduced diffusion is accompanied by an increase in dislocation density, a decrease in the formation of Kirkendall voids at the CIS/ GaAs heterojunction, and a roughening of the growth surface. All of these are consistent with a decreased atomic diffusion rate and probably a decreased vacancy concentration in the epitaxial layer as a result of the addition of Na. Further details will have to await the completion of the study. However, a reduced diffusivity would be consistent with the reduction or organization of native point defects in the material. 4. Conclusion Certainly, the jury is still out on the effect of Na. However, in spite of apparent contradictions in almost all of the experimental results, a general picture has emerged and the contradictions can be seen to be primarily the result of the experiments performed rather than necessarily any being specifically invalid. It is the conclusion of this author that the primary effect of Na is to organize the point defects during thin film growth of CIGS acting primarily as a surfactant without needing to enter the bulk of the material. As the improved surface structure is overgrown, the result is an improved bulk crystal quality. Furthermore, I conclude that Na is not necessary to device performances after growth is complete as long as it is available during an optimized growth process. Therefore, it is not surprising that there is no obvious correlation between Na at any particular site in the CIGS in completed devices and their performance. In support of this conclusion, I note that the opencircuit voltage, improved when adding Na, is correlated with the grains, not grain boundaries. If grain boundaries were the primary recombination point, one would expect grain size to have much more effect, especially when boundaries lie across the current collection pathway. All extended defects, including dislocations, second phases, twins, stacking faults, and grain boundaries, have been systematically excluded as having a clear influence on device performances [22,23]. This leaves only point defects as potential culprits. Furthermore, significant evidence for the existence of numerous defects in the energy gap of CIGS and their electronic effects has been found [16 21]. These defects are known to be affected by Na as summarized briefly above, although Na itself does not create new defects. At the same time, Na is widely thought based on device analyses to enhance carrier concentrations although it is not a dopant. Therefore, it must be increasing hole concentration by modifying the native defect concentration. This, in turn can affect the inplane conductivity of the films. The most likely target of Na is the Se vacancy or the In on Cu antisite defect. However, it may be sufficient that Na changes the tendency to organize these point defects into clusters, rendering them electrically inoffensive. Reduction in vacancy populations in particular would be consistent with observed reduced Ga diffusivity. Having provided an improved crystal structure during growth, Na is not needed in or on the grains in the final device, which is why a clear effect and sensitivity to residual Na levels is not observed. The conclusions are also consistent with the absence of an obvious physical location for Na activity or any systematic composition variation correlated with device performances. Acknowledgements I gratefully acknowledge the support of the U.S. Department of Energy through the National Renewable Energy Laboratory and the Basic Energy Sciences Program subcontract DEFG02-91ER Microanalysis was carried out in the Center for Microanalysis of Materials (CMM), supported by the Department of Energy. The author thanks all of the CMM staff for their invaluable help without which this research could not have been accomplished. 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