Stability of Cu(In,Ga)Se2 solar cells: a thermodynamic

Thin Solid Films 361±362 (2000) 338±345 www.elsevier.com/locate/tsf

Stability of Cu(In,Ga)Se2 solar cells: a thermodynamic approach J.F. Guillemoles Laboratoire d'Electrochimie et de Chimie Analytique (UMR 7575), ENSCP, 11 rue Pierre et Marie Curie, F 75231 Paris Cedex, France

Abstract Cu(In,Ga)Se2 (CIGS) based photovoltaic cells have demonstrated the highest solar energy conversion ef®ciencies ever for thin ®lm devices. They also exhibit excellent stability in ®eld tests and exceptional radiation hardness. The apparent paradox is that these results are obtained with a cell that contains a material that is chemically the most complex of the materials used in the various thin ®lm solar cells. Moreover, the device itself contains many elements, compounds and interfaces, all potential focus for evolution or reaction. Because of their central importance, the basic scienti®c foundations for the remarkable lifetime and stability of those devices are discussed, especially but not exclusively from a chemical point of view. A ®rst section is devoted to the assessment of the intrinsic thermodynamic stability of CIGS by a critical evaluation of available data. Its relationship with the formation energy of point defects is stressed. The chemical stability of the device interfaces are examined, including prospective buffer and window layers. q 2000 Elsevier Science S.A. All rights reserved. Keywords: CuInSe2; Cu(In,Ga)Se2; Solar cells; Stability; Thermochemistry

1. Introduction Cu(In,Ga)Se2 (CIGS) based photovoltaic cells have demonstrated the highest solar energy conversion ef®ciencies ever for thin ®lm devices, both for small area (18.8%) [1] and modules (12%) [2,3]. They also exhibit excellent stability in ®eld tests and exceptional radiation hardness. The success of CIGS in solar cells is a challenge to common sense. Good ohmic contact on p-type semiconductors is notoriously dif®cult, but it did not appear to be so between Mo and CIGS. Polycrystalline semiconductors tend to have reduced electronic performances, as compared to single crystals, due to carrier trapping or recombination at grain boundaries. This does not seem to be a problem in CIGS where grain boundaries are easily passivated. Semiconductors are generally very sensitive to impurities, those being generally lifetime killers. CIGS seems to be relatively immune to most impurities. Ironically, the one found to diffuse from the glass substrate (Na), was also found to improve the quality of the CIGS ®lms structurally and electronically. Moreover, the device itself contains many elements, compounds and interfaces, all potential focus for evolution or reaction. Last but not least, record ef®ciencies are obtained using a material that is chemically the most complex of the materials used in the various thin ®lm solar cells (a/Si, c-Si, CdTe,....), a compound that has also metastable states and shows signi®cant ionic conductivity. In the E-mail address: [email protected] (J.F. Guillemoles)

list above, one will recognise factors that plague or have plagued the development of other thin ®lm solar cells. Before discussing in details this apparent paradox, let us get an overview of the two questions addressed here: the intrinsic stability of CIGS and the global chemical stability of the device (for a detailed review, see [4], of which this work is a continuation, from where the elements of the overview were taken and where more references can be found). As far as terrestrial application are concerned, Cu(In,Ga)Se2-based solar modules have proven their stability in long term outdoor tests as well as under accelerated lifetime test conditions [5], and actually it is not uncommon that cells and modules show some improvement during testing [6], implying that the device as prepared is not optimized. This further suggests a positive evolution of the interfaces with time. The sensitivity of the cell to humidity either via the ZnO window [7] or via the bare CIGS absorber material [8] is mainly a concern at the production level: properly encapsulated, the modules are stable for years in spite of the chemical complexity of a system incorporating about ten different elements in layers of few tens of nm. More striking is the exceptional tolerance of CIGS to defects of various origins. Primo, one should keep in mind that CIGS is a non-stoichiometric compound, with deviations from stoichiometry in the % range: PV-grade material is generally obtained with a Cu content between 22 and 24%. Hence the main kind of defects are of the intrinsic nature, largely above the free carrier concentration. The material is strongly self-compensated, but in a way that

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J.F. Guillemoles / Thin Solid Films 361±362 (2000) 338±345

do not harm its electronic properties. This is quite surprising as generally, large deviations from stoichiometry are a problem for electronic applications as (e.g. in WOx, CuxSe, In2O32x, ZnO12x etc.) the valence electron unbalance is compensated by free carriers, so that relatively small deviations (in the 10 18 cm 23 range) already result in electronic degeneracy. In chalcopyrites though, electron unbalance from the 4 e 2 per atom rule which are larger by as much as three orders of magnitude, i.e. in the 10 21 cm 23 range) result in non-degenerate semiconductors. This implies that the formation of the compensating defects is energetically favoured as compared to formation of electronic carriers [9]. When this is the case, the Fermi level position is selfstabilized between a lower and upper value, as has indeed been found in many compound semiconductors [10,11]. Evaluation of the concentration of point defects in CIGS lead to the conclusion that many of the cationic defects have concentration above 10 18 cm 23 [4]. At those concentrations, they most likely form complexes, a fact also supported by ab-initio calculations [12]. Parallel to that ability to accommodate intrinsic defects, and certainly not unrelated, the electronic properties of the material appear less sensitive to impurities than is usually found in other semiconductors [4,13]. Secundo, we note the excellent radiation hardness of this type of solar cells, as compared to traditional space cells (Si, InP, InGaP) [14,15]. Other studies concluded that a very ef®cient, low temperature, defect recombination mechanism is needed to explain the exceptional radiation hardness [16,17]. Tertio, in spite of the high density of grain boundaries and other crystallographic defects, the transport properties of the free carrier did not seem to be affected, at least in the transversal direction. Accordingly, very high ef®ciencies could be obtained. An explanation for this feature was found involving the passivation of grain boundaries in air for the p-type material [18,19]. Surely, all this fortunate tolerance of the compound must be grounded in some of its particular properties. The most prominent one, from a solid state chemistry point of view, is its non-stoichiometry. Though the actual homogeneity domain of Cu(In,Ga)Se2 is not well known (see the discussion below), it is much larger than in the other semiconductors that have been successfully used for solar energy conversion (e.g. CdTe, Cu22xS). The existence of an electronic material with excellent electronic properties, compatible with high ef®ciency solar energy conversion, is quite puzzling in this context as most of the defects or defect complexes present must be inactive with respect to carrier recombination. Thus, their corresponding energy levels must either be shallow or altogether outside the bandgap. Again, ab-initio calculations on some defect complexes have shown that trend: defect levels of the considered complexes are be shallower than that of isolated point defects [12]. The second most important characteristic of CIGS is its mixed conductivity. Migration of Cu is well documented in

339

crystals and in thin ®lms [20±23] The room temperature Cu diffusion coef®cient in CuInSe2 (CIS) is in the range of 10 213±10 210 cm 2/s [24]. Regardless of the exact values, this implies that Cu can diffuse across the space charge region (SCR) of the CIGS absorber in a few minutes to several days at most [25]. Such a rapidly migrating species potentially could raise serious concerns as to the long-term stability of CIGS-based devices. Overall, these facts outline the picture of a compound with a `soft lattice' certainly related to the non-stoichiometry of the material and to its ionic conductivity. It is not surprising then that metastable electronic centers were found in this material. The most studied metastable level in CIGS is very different from those encountered for instance in a:Si on two aspects at least: they lead to an improvement of the device characteristics under operating conditions and they are reversible [26±28]. The `soft lattice' picture is central to the model proposed to account for the resilience and self-stabilization of CIGS [4,25,28]. It is because the material is highly disordered, but in a way that the most intrinsic defects complexes are not lifetime killers, and because Cu is mobile, that further defects (impurities, or radiation induced) can be self-passivated. This model uses the similarity between the point defects in a solid and solutes in an electrolyte, an analogy developed earlier [9], in that picture, the buffering action of the defect pool is essentially similar to the role of a buffer solution in aqueous electrolytes, where complexes in large concentrations (e.g. acids with a concentration such that the pH is close to their pKA) are used to ®x the ionic concentration, e.g. the pH. In other words, the defect complex pool acts as a reservoir, capable of receiving or delivering electronic and ionic charge carriers so that their net amount remains ®xed. Such defect reactions therefore act as both electronic and chemical buffers, by controlling both the electron concentration and the Cu concentration, respectively. In this model, a central role is given to mobile Cu as a vector of the buffering equilibrium. The detailed exposition of that mechanism being beyond the scope of this paper, the readers are referred to the original work [4,25]. This paper will ®rst elaborate further on the stability of CIGS from the thermodynamic point of view of CIGS and its connection with the interactions between point defects. Lastly, the stability in connection to the other partner compounds in the device will also be discussed, completing the results already presented [4]. 2. Intrinsic stability of CIGS Looking back at what has been achieved with CIGS, it is surprising to notice how little relatively is actually known about the compound. For instance, the extension of the homogeneity region has not been determined precisely, nor are the variations of most of the compound's properties with composition accurately known.

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2.1. Thermodynamic stability The homogeneity regions of CuInSe2 and CuGaSe2 at room temperature are not accurately known, but these two compounds were found miscible in all proportions though with a positive mixing enthalpy [29]. Most data report on the Cu2Se±In2Se3 or Cu2Se±Ga2Se3 tie-lines (here summarised as Cu2Se±III2Se3 tie-lines), which are relevant for most thin ®lm growth methods (because of the volatile Se and In2Se phases, the composition tends to adjust along the tie-line during the high temperature stages). In the range of temperatures used for thin ®lm growth (T , 7008C), each of these compounds was found to co-exist in equilibrium with Cu selenides on the III2Se3-poor side. The boundaries of the homogeneity region of the ternaries were also found to be in contact with ordered defect compounds (ODC) such as 1±3±5 or 1±5±8 ( CuIn3Se5 and CuGa5Se8, respectively) [30] on the III2Se3 rich side. Other known phase boundaries of the homogeneity region are Se [31,32], CuSe (but not InSe) and Cu2In4Se7 [33] and In [34]. Recent redetermination of the pseudo-binary phase diagram for CuInSe2 along the Cu2Se±In2Se3 tie line at T , 6008C, have shown [23,35] that the homogeneity range is narrower than what was observed before (24 to 24.5 at.% Cu) and would not include the nominal CuInSe2 composition (25 at.% Cu). The photovoltaically active ®lms of typical cell are prepared in presence of a large Se excess but result in ®lms of overall composition very close to the Cu2Se± III2Se3 tie-line. Their Cu content is generally between 22 and 24%, i.e. in the 1±3±5 1 CIGS two phase region. This reduces the risks of apparition of CuxSe secondary phases. In thin ®lms, the 1±3±5 phase is not observed by structural sensitive methods for Cu.22% [36], though the surface composition closely matches that of this compound for Cu,24.5% [37], while it is observed at lower Cu contents, so that its existence as a distinct surface phase on PVquality thin ®lms may be questioned [38]. Three factors apparently hinder the formation of 1±3±5 compound as a separate phase: (i) the addition of Ga and the presence of Na apparently inhibit the formation of the 135 phase [21] (substitution of part of the In atoms in the CIS lattice by Ga atoms, or growth in assisted with Na migration from the glass substrate, led to an increase of the homogeneity domain of the chalcopyrite phase along the pseudo-binary tie-line in the (In,Ga)2Se3 direction, this effect of Ga is further substantiated by ab-initio calculation [39]), (ii) the ®lms are prepared at a temperature where the CIGS phase has a larger domain of existence and then cooled down to room temperature [35] so that the solid solution may be frozen-in, (iii) the nucleation of a weakly n-type phase (1± 3±5) in a p-type phase (CIGS) is inhibited by an important interfacial energy of electrostatic origin [38,40] and by lattice mismatch [36]. Because CIS is the base compound for solar cells, because the system was more investigated and for simplicity

as CIS retains the main features of the CIGS quaternary, we will conduct the discussion on this ternary with only mention of the effect of Ga addition. The fact that tie-lines cannot cross combined with the system properties exposed above, determines most of the tie-lines of the system. These are reported as a continuous line on Fig. 1. Besides those that have strong experimental support, other tie-lines are more problematic to determine. A ®rst instance is the existence of the Cu±CuInSe2 tie line or alternatively of the Cu2Se±Cu2In. Experimental data on ®lms prepared from metallic precursors do not enable a clear choice. The second instance is the CuInSe2±In-In4Se3 triangle. There, the only constraint is that InSe±CIS is not a tie line, hence Cu2In4Se7 has a tie line either with In4Se3 or with In. On ambiguous cases, the choices made (dotted lines) are based on compatibility with free energies of formations, as we turn to discuss now. 2.2. Free energy values and point defects Most of the free energies of formation of the established compounds in the Cu±In±Se system are known. They can be obtained from various sources [41±47]. Those on the Cu±Se binaries show little discrepancies and are considered reliable. On the In±Se binary more con¯icting values are found. These have been discussed in [46]. Some values are available on the Cu±In binary [48], but the phase diagrams published show considerable differences. Fortunately, these were shown to be of little consequence for what follows as the free energies of formation of the metallic alloys are rather small compared to the metal±non metal ones. In those systems, we have ®rst used the fact that the free energy per atom must be a convex function of molar fraction [49] to eliminate or adjust incorrectly determined values using the best agreed ones as a guide. The knowledge

Fig. 1. Phase diagram used in this study. Dotted lines correspond to assumed tie-lines.

J.F. Guillemoles / Thin Solid Films 361±362 (2000) 338±345 Table 1 Standard free energies (298 K) of some compounds of the Mo/CIGS/CdS/ ZnO system. The enthalpies of formations of the elements at 258C, 1 bar and their entropies at 0 K are set to zero. Compound G8 (kJ/mole) Ref.

Compound G8 (kJ/mole) Ref.

Cu8 In8 Se8 Cu2Se Cu7Se4 Cu3Se2 CuSe CuSe2 In2Se3 InSe In3Se4 Cu11In9 Cu2In CuInSe2 Cu2In7Se4 CuIn3Se5 CuIn5Se8 Ga2Se3 Ga8

CdO CdS CdSe Cu2O Cu2S CuO CuS In2O3 In2S3 MoSe2 Mo8 ZnO ZnS ZnSe SnO2 SnS2 SnSe2 Ga2O3 Ga2S3

29.9 217.2 212.6 2103.9 2351 2173 265.2 281 2360 2130 2416.5 2410 261.5 2280 2942.5 2661.5 21041.6 2462.4 212.17

[41] [41] [41] [41] Calculated Calculated [41] [48] [65] [46] [46] [43] [48] Calculated Calculated Calculated Calculated [41] [41]

2274.7 2176.8 2170.7 2198.2 2115.6 2168.8 273.6 2957.0 2414.4 2180.2 28.5 2363.6 2222.2 2193.5 2596.43 2179.62 2154.89 21114.4 2558.7

[41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41] [41]

of the established tie-lines of the system was also used to assess the plausibility of the free energies of formations available and to propose a plausible range for those. For instance, the existence of a CIS±In tie line instead of a In2Se3 (or InSe) ±Cu11In9 tie line imposes a lower bound to the formation energy of CIS as discussed in more details below. The values used in this study not directly taken from reference tables because of internal inconsistencies were interpolated by assuming a smooth (and convex) variation of the free energy per atom with composition. The remaining problem is to get a reliable value of the free energy of formation of CIGS and of the ordered compounds found on the Cu2Se±In2Se3 tie line. Experimental values are available for CuInSe2 [41,50,51], showing the ternary is some 10±80 kJ more stable than the mixture of binaries Cu2Se 1 In2Se3 [43,45], with a trend that experimentally determined values are signi®cantly smaller than those expected from various models. Ab-initio (computed) values enable to calculate (using the hexagonal In2Se3 free energy of formation given in [10] and speci®c heat values of [52]) a stabilisation free energy of 52 kJ for CIS. A minimal stabilisation energy of about 50 kJ/M for CIS was found necessary to get the CIS±In tie line with the choice of formation energies calculated to determine the free energies of formation of the Cu2In4Se7, CuIn3Se5 and CuIn5Se8 compounds. We have used (or interpolated in the case of Cu2In4Se7) the 0 K formation energies from CIS and In2Se3 of [12] and interpolated the speci®c heat values reported on the Cu2Se±In2Se3 tie line. The obtained values have been then used with the calculated values of CIS and In2Se3 to yield the free energies of formation of the ODC compounds. No estimates have been made for CIGS. The stability of CuGaSe2 versus its constituent binaries is postulated to

341

similar as that of CIS. A fact supported by the similarities of the Cu±Ga±Se and Cu±In±Se systems. The two compounds are miscible in all proportions but their mixture has a positive heat of mixing that can make the alloy up to 3.3 kJ/M less stable than the ideal mixture [29]. Overall, the values considered in this study are presented in Table 1. Using Table 1 and Fig. 1, it is now a straightforward matter to compute the chemical potential of Cu, In and Se at the phase boundaries. The data is reported on Fig. 2. Those values are important in connection of the stability of the material with respect to out-diffusion, mainly of Cu, and they will be used in the next section. But let us ®rst discuss them in connection to the intrinsic stability and the electrical stability of that compound. The dependence on chemical potential of the point defect free energy of formation is well known. As CIGS thin ®lms for PV application are grown in presence of excess In and Se, they are located at the bottom left of the existence domain of Fig. 2, i.e. with mCu between 20.4 and 20.7 eV and mIn between 21 and 21.6 eV, which reduces considerably the free energy of formation of the related vacancies. It is relatively simple to measure the chemical potential (m) of Cu in CIS or CIGS using an electrochemical set-up [53,54], and the measured values correlate quite well with the present calculated ones, if one considers the possible inaccuracy of the of the thermodynamic values (the error is estimated to be below 0.2 eV). Measurements gave mCu of 20.3 to 20.4

Fig. 2. Stability diagram of CuInSe2 with respect to its neighbouring phases. The arrows indicate which phase is approached when (1) VCu, (2) VIn (3), VSe (4) {InCu, 2VCu} point defects are created.

342


Table 2 Lower boundaries for some defect free energies of formation (energies in eV). VCu

VIn

VSe

InCu

CuIn

SeIn

InCu 1 2VCu

CuIn 1 2Cui

VSe 1 CuIn

SeCu

InSe

CuSe

0.75

1.88

2.21

0.22

1.47

1.88

0.65

0.89

2.35

0.75

2.21

2.1

eV for crystals grown from a stoichiometric mixture, while crystals grown from a 2% Se rich melt had somewhat lower values, between 20.45 and 20.5 eV [54]. Recent measurements on PV-grade CIGS thin ®lms yielded mCu of 20.55 eV [55], also in agreement with the present analysis. The limiting values of mCu inferred from Fig. 2 are very close to the calculated [12] as well as experimentally obtained [53,54] standard free energies of formation of VCu. Note that the standard free energy of formation of a point defect is referred to the standard state of the elements. For instance in the reaction of cation vacancy creation, the neutral cation, i.e. the element, removed from the lattice is delivered into a reservoir where its chemical potential is zero (that of the element in its standard state) [12]. It is contended here that this closeness is no coincidence but comes from the fact that the free energy of formation of a defect cannot become negative (in that case the compound would be unstable). Indeed, if a compound had a free energy of reaction of formation of a point defect that is negative or null, according to the precise de®nition of that energy given above, that defect would form spontaneously, i.e. the total free energy of the compound would not be minimal. This is another way of stating that the compound would not have attained its thermodynamic equilibrium, in contradiction with the premises. This is a very useful statement once it is realised that the limiting chemical potentials for the CIS phase stability calculated here give us a lower bound for the standard free energy of formation of the related defect. For instance, since the calculated chemical potential of Cu is higher than 20.75 (^0.2) eV here, the standard free energy of formation of VCu is higher than 0.75 (^0.2) eV because the free energy of the reaction CuCIS ! Cu8 1 VCu leading to spontaneous formation of vacancies should be positive. Indeed, using ab-initio values [12], the chemical potential of Cu in CIS can be deduced in the triangle CIS±Cu2In4Se7± Se, i.e. where it is the most negative. There mCu 20:335 eV is found. A value higher than the opposite of the standard free energy of formation computed for VCu in the same work (0.6 eV). Similarly, a lower boundary for the standard free energies of formation of other point defects can be deduced (Table 2). It is clear from Figs. 1 and 2 that the values of the chemical potential where they would form the most easily corresponds to the contact with the phase that is formed when the point defect is created: the creation of CuIn in CIS ultimately leads to the formation of CuSe. This again is certainly no coincidence. A non-orthodox but yet illuminating view on a phase transition is that of point defects that form and cluster as composition (or temperature) changes

and eventually form a nuclei for a new phase. In that view, the defect or defect complexes with the lowest energy of formation are those that make the phase unstable upon ordering when the composition shifts. A splendid example of that is given in [12] where it is described how the formation of {InCu; 2VCu} complexes are the precursors for the next ODC phases. It is no surprise by the way that the lower bound of the value for the formation of such complexes found by the phase stability rule equals that given in [12], i.e. the stability of the 2±4±7 phase was found using their values. Hence the lower bounds given in Table 3 can be considered as good approximations of the defect formation energies for those defects that are the cause of the phase transition. {InCu; 2VCu}complexes and VCu are apparently such defects. The difference between the standard free energy of formation and that lower bound being by de®nition equal to the ordering energy, the latter is certainly small Table 3 Free energies of reaction for some of the relevant reactions at the interfaces in the Mo/CIGS/buffer/TCO type of solar cells. Reactions Back contact Mo 1 2Cu2Se ) MoSe2 1 4Cu 3Mo 1 2In2Se3 ) 3MoSe2 1 4In 3Mo 1 2Ga2Se3 ) 3MoSe2 1 4Ga Buffer and Window Cu2Se 1 CdS ! Cu2S 1 CdSe In2Se3 1 (Zn,Cd)Cd(S,Se) ! (Zn,Cd)In2(S,Se)4 In2Se3 1 3CdS ! 3CdSe 1 In2S3 Ga2Se3 1 3CdS ! 3CdSe 1 Ga2S3 Ga2Se3 1 3ZnS ! 3ZnSe 1 Ga2S3 CdS 1 ZnO ! CdO 1 ZnS Cu2Se 1 ZnO ! Cu2O 1 ZnSe In2Se3 1 3ZnO ! In2O3 1 3ZnSe Ga2Se3 1 3ZnO ! Ga2O3 1 3ZnSe 2Ga2Se3 1 3SnO2 ! 2Ga2O3 1 3SnSe2 2In2Se3 1 3SnO2 ! 2In2O3 1 3SnSe2 2Cu2Se 1 SnO2 ! Cu2O 1 SnSe2 In2S3 1 3ZnO ! In2O3 1 3ZnS

DG (kJ)

Possible

23.5 1188

Yes No

1361

No

25.6 ,0

Yes Yes

0.3

?

278

Yes

210.2

Yes

43.5 75.8 260.3

No No Yes

2141.7

Yes

20.62

No

183.4

No

252.9

No

2128.4

Yes


343

for many possible defects of the compound. This is likely to be traced back to the multinary nature of CIGS: the number of point defects grows exponentially with the number of constituents thus increasing the number of ways for the system to minimize its free energy. While in simple systems, e.g. binaries, this is usually achieved by ordering, it is less necessary so when the complexity of the system increases, giving it more ¯exibility to accommodate defects.

reaction 2 and 3). With the CdS interface, a high content of Ga will tend to destabilise the CIGS versus the decomposition reaction products as the energy of reaction is large. The additional problem here rests with the existence of de®ned compounds and solids solutions in the Cu(In,Ga)(S,Se)2±(Zn,Cd)(S,Se) system [61] (reactions 4 and 5). The free energy of formation provides an additional thermodynamic driving force for intermixing. We can therefore expect some intermixing to occur at this interface.

3. Interfacial stability

Cu2 Se 1 CdS ! Cu2 S 1 CdSe

2

We now turn to the discussion of the stability of the full device, i.e. of the stability of CIGS in conjunction with the other compounds that are present in the device. Values of the standard free energies of formation of the relevant materials were collected from various sources [41±48] and are given in Table 1. In this section, instead of making an hypothesis on the exact value of the stabilisation energy, we have considered reactions with the constituent binaries. For those reactions that are allowed with the binaries, the thermodynamic driving force has to be discussed in connection with the stabilisation energy. As guidelines, the values reported above will be used in the discussion when needed to evaluate the formation of CIGS as

2CuInSe2 1CdS ! Cu2 S 1 CdSe 1 In2 Se3

3

xCuInSe2 1CdSe ! Cux CdInx Se2x11

4

CdSe 1 In2 Se3 ! CdIn2 Se4

5

1=2f Cu2 Se1xIn2 Se3 112xGa2 Se3 g ! CuInx Ga12x Se2 1 3.1. Mo/Cu(In,Ga)Se2 None of the studies conducted on the Mo/CIGS interface ever found evidence of Mo diffusion into the CIGS ®lm [56,57]. A characteristic feature of that interface is the formation during growth of an interfacial MoSe2 layer , 10±100 nm thick with the help of the excess Se which is used in all methods leading to a photovoltaic-quality ®lm [58±60]. Using Table 3, it is readily concluded that the MoSe2-forming reaction with each of the three binary compounds is thermodynamically forbidden, with the exception of Cu2Se. But in this case, the free energy gain of the reaction (about 1 kJ/M Cu) is less than the stabilization energy, even considering the lowest value of CIS. Thus, in the absence of excess Se, which is found during growth only, the reaction with CIGS cannot proceed. Some compounds have been reported in the JCPDS tables in the In±Mo±Se, e.g. InMo3Se6, InxMo15Se19, InMo3Se3, and Ga±Mo±Se, e.g. GaMo4Se8 systems. Though their energies of formation are not known, their presence can be discarded as they are expected to be in Se-poor ambiants only, i.e. not the conditions leading to PV quality material. 3.2. Cu(In,Ga)Se2/buffer stability According to Table 3 reactions of the binaries with CdS are possible, but except for Ga, the energies involved are smaller than the lower stabilisation energy for CIS (e.g.

This was somewhat con®rmed recently by PES experiments on the interface [62]. The stability of this interface is also kinetic, and limited by the ability of the constituents (Cu, Ga, Cd, In, Se, S) to diffuse. Except Cu, none of those has a high diffusion coef®cient at room temperature, so that only limited (nm range) interdiffusion is expected. Some reaction of CdS was observed with copper-rich CIS and with CGS. These observations can be rationalised with the present thermodynamic calculations: Cu-rich CIS contains some CuxSe that is probably the cause of the observed reaction, whereas for CGS effective decomposition of the compound is possible. A more detailed discussion is given in [4]. The case of ZnS shows the same trend, but is found much more stable (Table 3). This buffer might then be preferred on high Ga content material (high band gap cells). In2(S,Se)3 are also possible buffers. They are not thermodynamically stable on CIGS, as can be inferred from the phase diagram (Fig. 1), and upon intermixing they will form the sequence of ODC. As with the buffers considered above, the possible use of these buffers comes from the slow interdiffusion at room temperature. The actual long-term stability of these interface is an open question. 3.3. Buffer/TCO stability The direct deposition of ZnO on CIGS leads to the formation of In and Ga oxides (Table 3). Those are not necessarily harmful to the junction by themselves, but they lead to a change in composition of the absorber surface (enrichment in Cu). The effect is much less pronounced with ZnS than with ZnO, so that the former could also be used as a buffer layer. The buffer layers based on In or Ga sulphides or selenides could be ef®ciently oxidised during the ZnO deposition, which may question their use. We note again the strong tendency of Ga compounds to react. This reactivity of Ga may be one of the reason why realisation of high band gap cells actually leads to severe voltage losses. A stable and well controlled interface will probably require buffer layers that will enable to limit and control the surface

344


oxidation of CIGS. Alternatively we note that SnO2-based window are much less reactive and could be safely used as TCO. Since Sn is also a n-type doping agent in CIS, there is the hope that a limited interdiffusion may lead to the surface electrical inversion as was found with Cd or Zn compounds. 3.4. Cu out-diffusion We address now the speci®c problem posed by the relatively fast Cu migration in CIGS, as well as in most materials. We have looked at the thermodynamic driving forces for Cu migration by measuring the Cu chemical potential in various cell-relevant materials. The measured value for CIGS was already given above (20.5 V/Cu8). Using different cells but following the same principle, we have measured the (preliminary) values in CdS CBD (20.4 V/ Cu8) and in MoSe2 (20.2 V/Cu8) [63]. They were found to change steeply with Cu incorporation, hence the range given in Fig. 3. The later value for CdS is somewhat larger than the one deduced from solubility experiments in CdS [64]. Considering additionally to the chemical forces, the electrostatic forces (the built-in ®eld), the pro®le of the electrochemical potential in the cell shows that Cu in CIGS is in a potential well, with no driving force to escape. Actually, out-diffusion in CdS were only found for Cu rich material, where the chemical potential of Cu is much closer to Cu8 as is also easily deduced from Fig. 2. In summary, whereas Cu is mobile in CIGS, thanks to the low chemical potential it has in this compound and to the direction of the electric ®eld, it is ef®ciently trapped an is harmless to the neighbouring compounds as well as immune to precipitation (over-potential of several hundreds of mV would need to be generated, a situation unlikely to arise with mere ohmic drops). The situation is thus quite different as the one found in Cu2S solar cells, as here Cu is more a good fortune for its role in the self healing process than a severe drawback [4].

4. Conclusion In short, what makes CIGS so special among semiconductors used for opto-electronic applications? The thermodynamics and the solid-state chemistry of the system provide some answers: the unusual stability of CIGS comes from its ability to accommodate defects chemically and electrically (radiation-induced, impurities, dislocations) rather than from the dif®culty to actually form defects. In that sense, CIGS is elastic rather than plastic, ¯exible rather than hard. This endows a new type of resiliency to the material, the kind of which that constitutes a lesson for future material engineering [25]. The point of view put forward here is that the puzzling properties of CIGS are tied together, i.e. they originate in the same chemical properties that are also the roots of the non-stoichiometry and of the ionic conduction of the compound, and of their interactions that result in the previously described self-healing behaviour [4,25]. These properties can be traced both to the microscopic point of view (the accommodation of intrinsic point defect) and to the macroscopic ones (the stability diagram), and we have seen how they are connected, with the result that macroscopic thermodynamic data were shown to be an indicator on the formation of point defects when the ordering energy is small, i.e. when the system is able to accommodate nonstoichiometry. From the device point of view, we see that the system actually used is already quite good chemically. It is possible to change the CIGS partner in the device e.g. to ZnS or SnO2, for an improved stability, but electrical parameters, such as band alignment, have also to be taken into account. The existing buffer seems to be the best compromise to date between chemical and electrical requirements. The possible interdiffusion between CdS and CIGS is possibly bene®cial, when limited to a short range (few nm), in limiting the interface recombination and preparing a buried junction. The not uncommonly observed improvement of cell performances with ageing could ®nd its origin here. Last but not least, it was found that Ga addition in CIGS resulted in less stable interfaces, which may be one of the reason of the dif®culties of preparing optimal high band gap cells in the system.

Acknowledgements

Fig. 3. Representation of the range of values of mCu, the copper chemical potential, in the various phases present in a solar cell. The actual driving force for migration, i.e. h Cu, the electrochemical potential of Cu, is represented below.

This work originates in the joint effort carried on by L. Kronik and D. Cahen (Weizmann Inst.), R. Herberholtz, U. Rau and H.W. Schock (IPE), and the author, where seminal idea where ®rst formulated. It also owes much to stimulating discussions with J. Vedel (ENSCP). It was carried out with the support of the EC under Joule contract (#JOR3-CT970149). The author is grateful to S. Cassaignon and E. Clolus for the electrochemical measurements.


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