With a power conversion efficiency of 18.8% on a 0.5-cm² laboratory cell [1] and 16.6% for mini-modules with an area of around 20 cm² [2], Cu(In,Ga)Se₂ is established as an efficient thin-film solar cell technology. The start of production at several places provides a new challenge for research on this material. However, these recent achievements are based on a long history of research and technological development.

CuInSe₂ was synthesised for the first time by Hahn in 1953 [3]. In 1974, this material was proposed as a photovoltaic material [4] with a power conversion efficiency of 12% for a single-crystal solar cell. In the years 1983–84, Boeing Corp. reported efficiencies in excess of 10% from thin poly crystalline films obtained from a three-source co-evaporation process [5]. In 1987 Arco Solar achieved a long-standing record efficiency for a thin-film cell of 14.1% [6]. It took a further ten years, before Arco Solar, at that time Siemens Solar Industries (now Shell Solar), entered the stage of production. In 1998, the first commercial Cu(In,Ga)Se₂ solar modules were available [7]. In parallel, a process which avoids the use of H₂Se is being developed by Shell Solar in Germany [8] (see also Chapter IV-1). Other companies in the USA, Global Solar and ISET, plan to commercialise modules prepared on other than glass substrates. In Europe, the long-term development efforts of the EUROCIS consortium on the co-evaporation process resulted in the activity of Würth Solar with pilot production envisaged in 2003 [9,10]. In Japan, two lines for film preparation are planned by Showa Shell (selenisation by H₂Se) [11] and Matshushita (co-evaporation) [12].

In this chapter, we give a short overview on the present knowledge of Cu(In,Ga)Se₂-based heterojunction thin film solar cells. We focus on four points: (i) The description of the basic material properties such as crystal properties, phase diagram, and defect physics. (ii) Description of the cell technology starting from the growth of the polycrystalline Cu(In,Ga)Se₂ absorber up to device finishing by heterojunction formation and window layer deposition. This section also discusses basic technologies for module production. (iii) The electronic properties of the finished heterostructure. (iv) Finally, Section 5 discusses the photovoltaic potential of wide-gap chalcopyrites, namely CuGaSe₂ and CuInS₂, as well as that of the pentenary alloy system Cu(In,Ga)(S,Se)₂ and the possibility of building graded-gap structures with these alloys.

This chapter can only briefly cover those scientific issues that are relevant for photovoltaic applications. More detailed information can be found in two review articles by the present authors [13,14] as well as in references [15–19].

2 Material Properties

2.1 Chalcopyrite Lattice

CuInSe₂ and CuGaSe₂, the materials that form the alloy Cu(In,Ga)Se₂, belong to the semiconducting I–III–VI₂ materials family that crystallise in the tetragonal chalcopyrite structure. The chalcopyrite structure of, for example, CuInSe₂ is obtained from the cubic zinc blende structure of II–VI materials like ZnSe by occupying the Zn sites alternately with Cu and In atoms. Figure 1 compares the two unit cells of the cubic zinc blend structure with the chalcopyrite unit cell. Each I (Cu) or III (In) atom has four bonds to the VI atom (Se). In turn each Se atom has two bonds to Cu and two to In. Because the strengths of the I–VI and III–VI bonds are in general different, the ratio of the lattice constants c/a is not exactly two. Instead, the quantity 2–c/a (which is −0.01 in CuInSe₂, +0.04 in CuGaSe₂) is a measure of the tetragonal distortion in chalcopyrite materials.

FIGURE 1 Unit cells of chalcogenide compounds. (a) Sphalerite or zinc blende structure of ZnSe (two unit cells); (b) chalcopyrite structure of CuInSe2. The metal sites in the two unit cells of the sphalerite structure of ZnSe are alternately occupied by Cu and In in the chalcopyrite structure.

2.2 Band-Gap Energies

The system of copper chalcopyrites Cu(In,Ga,Al)(Se,S)₂ includes a wide range of band-gap energies E_g from 1.04 eV in CuInSe₂ up to 2.4 eV in CuGaS₂. and even 2.7 eV in CuAIS₂, thus, covering most of the visible spectrum. All these compounds have a direct band gap, making them suitable for thin film photovoltaic absorber materials. Figure 2 summarises lattice constants a and band-gap energies E_g of this system. Any desired alloys between these compounds can be produced as no miscibility gap occurs in the entire system. We will discuss the status and prospects of this system in Section 5 in more detail.

FIGURE 2 Band-gap energies E_g versus the lattice constant a of the Cu(In,Ga,Al)(S,Se)₂ alloy system.

2.3 The Phase Diagram

Compared with all other materials used for thin-film photovoltaics, Cu(In,Ga)Se₂ has by far the most complicated phase diagram. Figure 3 shows the phase diagram of CuInSe₂ given by Haalboom et al. [20]. This investigation had a special focus on temperatures and compositions relevant for the preparation of thin-films. The phase diagram in Figure 3 shows the four different phases which have been found to be relevant in this range: the α-phase (CuInSe₂), the β-phase (CuIn₃Se₅), the δ-phase (the high-temperature sphalerite phase) and CuySe. An interesting point is that all neighbouring phases to the α-phase have a similar structure. The β-phase is actually a defect chalcopyrite phase built by ordered arrays of defect pairs (Cu vacancies V_Cu and In–Cu antisites In_Cu). Similarly, CuySe can be viewed as constructed from the chalcopyrite by using Cu–In antisites Cu_ln and Cu interstitials Cu_i. The transition to the sphalerite phase arises from disordering the cation (Cu, In) sub-lattice, and leads back to the zinc blende structure (cf. Figure 1 (a)).

FIGURE 3 Quasi-binary phase diagram of CuInSe₂ along the tie-line that connects the binary compounds In₂Se₃ and Cu₂Se established by Differential Thermal Analysis (DTA) and microscopic phase analysis (After Haalboometal. [20]). Note that at 25 at.% Cu no single phase exists.

The existence range of the α-phase in pure CuInSe₂ on the quasi-binary tie line Cu₂Se-In₂Se₃ extends from a Cu content of 24 to 24.5 at.%. Thus, the existence range of single-phase CuInSe₂ is relatively small and does not even include the stoichiometric composition of 25 at.% Cu. The Cu content of absorbers for thin-film solar cells varies typically between 22 and 24 at.% Cu. At the growth temperature this compositional range lies within the single-phase region of the α-phase. However, at room temperature it lies in the two-phase α+β region of the equilibrium phase diagram [20]. Hence one would expect a tendency for phase separation in photovoltaic-grade CuInSe₂ after deposition. Fortunately, it turns out that partial replacement of In with Ga, as well as the use of Na-containing substrates, considerably widens the single-phase region in terms of (In+Ga)/(In+Ga+Cu) ratios [21]. Thus, the phase diagram hints at the substantial improvements actually achieved in recent years by the use of Na-containing substrates, as well as by the use of Cu(In,Ga)Se₂ alloys.

2.4 Defect Physics of Cu(In,Ga)Se₂

The defect structure of the ternary compounds CuInSe₂, CuGaSe₂, CuInS₂, and their alloys, is of special importance because of the large number of possible intrinsic defects and the role of deep recombination centres for the performance of the solar cells. The features that are somewhat special to the Cu-chalcopyrite compounds are the ability to dope these compounds with native defects, their tolerance to large off-stoichiometries, and the electrically neutral nature of structural defects in these materials. It is obvious that the explanation of these effects significantly contributes to the explanation of the photovoltaic performance of these compounds. Doping of CuInSe₂ is controlled by intrinsic defects. Samples with p-type conductivity are grown if the material is Cu-poor and annealed under high Se vapour pressure, whereas Cu-rich material with Se deficiency tends to be n-type [22,23]. Thus, the Se vacancy V_Se is considered to be the dominant donor in n-type material (and also the compensating donor in p-type material), and the Cu vacancy V_Cu the dominant acceptor in Cu-poor p-type material.

By calculating the metal-related defects in CuInSe₂ and CuGaSe₂, Zhang et al. [24] found that the defect formation energies for some intrinsic defects are so low that they can be heavily influenced by the chemical potential of the components (i.e., by the composition of the material) as well as by the electrochemical potential of the electrons. For V_Cu in Cu-poor and stoichiometric material, a negative formation energy is even calculated. This would imply the spontaneous formation of large numbers of these defects under equilibrium conditions. Low (but positive) formation energies are also found for the antisite Cu_In in Cu-rich material (this defect is a shallow acceptor which could be responsible for the p-type conductivity of Cu-rich, non-Se-deficient CuInSe₂). The dependence of the defect formation energies on the electron Fermi level could explain the strong tendency of CuInSe₂ to self-compensation and the difficulties of achieving extrinsic doping. The results of Zhang et al. [24] provide a good theoretical model of defect formation energies and defect transition energies, which exhibits good agreement with experimentally obtained data. Table 1 summarises the ionisation energies and the defect formation energies of the 12 intrinsic defects in CuInSe₂. The energies (bold values in Table 1) for V_Cu. V_In, Cu_I, Cu_In, In_Cu, are obtained from a first principle calculation [24] whereas the formation energies in italics (Table 1) and for the other defects are calculated from the macroscopic cavity model [25]. The ionisation energies used in Table 1 are either taken from Zhang et al. [24] or from the data compiled in reference [26]. Note that the data given in references [25,26] for the cation defects differ significantly from those computed in reference [24].

TABLE 1 Electronic Transition Energies and Formation Energies ΔU of the 12 Intrinsic Defects in CuInSe₂

Source: the Ionisation Energies in Italics are Derived from Reference [26], and the Formation Energies in Italics are from Reference [25]. All the Numbers in Bold Type are from Reference [24].

^aDifference between the valence/conduction band energy for acceptor/donor states.

^bFormation energy ΔU of the neutral defect in the stoichiometric material.

^cCovalent.

^dIonic.

Further important results in reference [24] are the formation energies of defect complexes such as (2V_Cu, In_Cu). (Cu_In, In_Cu) and (2Cu_I, Cu_In), where Cu_I is an interstitial Cu atom. These formation energies are even lower than those of the corresponding isolated defects. Interestingly, the (2V_Cu, In_Cu) complex does not exhibit an electronic transition within the forbidden gap, in contrast to the isolated In_Cu-anti-site, which is a deep recombination centre. As the (2V_Cu, In_Cu) complex is most likely to occur in In-rich material, it can accommodate a large amount of excess In (or likewise deficient Cu) and, at same time, maintain the electrical performance of the material. Furthermore, ordered arrays of this complex can be thought as the building blocks of a series of Cu-poor Cu–In–Se compounds such as CuIn₅Se₅ and CuIn₅Se₈ [24].

Let us now concentrate on the defects experimentally detected in photovoltaic grade (and thus In-rich) polycrystalline films. In-rich Cu(In,Ga)Se₂ is in general highly compensated, with a net acceptor concentration of the order of 10¹⁶ cm⁻³. The shallow acceptor level V_Cu (which lies about 30 meV above the valence band) is assumed to be the main dopant in this material. As compensating donors, the Se-vacancy V_Se as well as the double donor In_Cu are considered. The most prominent defect is an acceptor level at about 270–300 meV above the valence band, which is reported by several groups from deep-level transient spectroscopy [27] and admittance spectroscopy [28,29]. This defect is also present in single crystals [30]. The importance of this transition results from the fact that its concentration is related to the open-circuit voltage of the device [31–33]. Upon investigating defect energies in the entire Cu(In,Ga)(Se,S)₂ alloy system, Turcu et al. [34] found that the energy distance between this defect and the valence band maximum remains constant when alloying CuInSe₂ with Ga, whereas the energy distance increases under S/Se alloying. Assuming that the defect energy is independent from the energy position of the band edges, like the defect energies of transition metal impurities in III–V and II–IV semiconductor alloys [35,36], the authors of reference [34] extrapolate the valence band offsets ΔE_v = −0.23 eV for the combination CuInSe₂/CuInS₂ and ΔE_v = 0.04 eV for CuInSe₂/CuGaSe₂. Recently, transient photocapacitance studies by Heath et al. [37] unveiled an additional defect state in Cu(In,Ga)Se₂ at about 0.8 eV from the valence band. Again, the defect energy is independent of the Ga content in the alloy. Figure 4 summarises the energy positions of bulk defects in the Cu(In₁₋_xGa_x)Se₂ and the CuIn(Se₁₋_yS_y)₂ alloy system with the defect energy of the bulk acceptor used as a reference energy to align the valence and conduction band energy [38]. Additionally, Figure 4 shows the activation energy of an interface donor. This energy corresponds to the energy difference ΔE_Fn between the Fermi energy and the conduction band minimum at the buffer absorber interface [39]. Notably this energy difference remains small upon alloying CuInSe₂ with Ga, whereas ΔE_Fn increases when alloying with S.

FIGURE 4 Band gap evolution diagram of the CuIn(Se,S)₂ (a) and the Cu(In,Ga)Se₂ (b) alloy system with the trap energy E_T(N₂, open diamonds) taken as an internal reference to align the conduction band and the valence band energies E_c and E_v. The energy position of an additional defect state in Cu(In,Ga)Se₂ (full diamonds) as well as that of an interface donor (open triangles) in Cu(In,Ga)(Se,S)₂ is also indicated.

3 Cell and Module Technology

3.1 Structure of the Heterojunction Solar Cell

The complete layer sequence of a ZnO/CdS/Cu(In,Ga)Se₂ heterojunction solar cell is shown in Figure 5. The device consists of a typically 1-μm-thick Mo layer deposited on a soda-lime glass substrate and serving as the back contact for the solar cell. The Cu(In,Ga)Se₂ is deposited on top of the Mo back electrode as the photovoltaic absorber material. This absorber layer has a thickness of 1–2 μm. The heterojunction is then completed by chemical bath deposition (CBD) of CdS (typically 50 nm thick) and by the sputter deposition of a nominally undoped (intrinsic) i-ZnO layer (usually of thickness 50–70 nm) and then a heavily doped ZnO layer. As ZnO has a band-gap energy of 3.2 eV it is transparent for the main part of the solar spectrum and therefore is denoted as the window layer of the solar cell.

FIGURE 5 Schematic layer sequence of a standard ZnO/CdS/Cu(In,Ga)Se₂ thin-film solar cell.

3.2 Key Elements for High-Efficiency Cu(In,Ga)Se₂ Solar Cells

We first mention four important technological innovations which, during the decade 1990–2000, have led to a considerable improvement of the efficiencies and finally to the record efficiency of 18.8% [1]. These steps are the key elements of the present Cu(In,Ga)Se₂ technology:

• The film quality has been substantially improved by the crystallisation mechanism induced by the presence of Cu_ySe (y < 2). This process is further supported by a substrate temperature close to the softening point of the glass substrate [40].

• The glass substrate has been changed from Na-free glass to Na-containing soda-lime glass [40,41]. The incorporation of Na, either from the glass substrate or from Na-containing precursors, has led to an enormous improvement of the efficiency and reliability of the solar cells, as well as to a larger process tolerance.

• Early absorbers consisted of pure CuInSe₂. The partial replacement of In with Ga [42] is a further noticeable improvement, which has increased the band gap of the absorber from 1.04 eV to 1.1–1.2 eV for the high-efficiency devices. The benefit of 20–30% Ga incorporation stems not only from the better band gap match to the solar spectrum but also from the improved electronic quality of Cu(In,Ga)Se₂ with respect to pure CuInSe₂ [21,32].

• The counter electrode for the CuInSe₂ absorber of the earlier cells was a 2 μm thick CdS layer deposited by Physical Vapour Deposition (PVD). This has been replaced by a combination of a 50-nm-thin CdS buffer layer laid down by chemical bath deposition [43,44] and a highly conductive ZnO window layer.

3.3 Absorber Preparation Techniques

3.3.1 Basics

The preparation of Cu(In,Ga)Se₂-based solar cells starts with the deposition of the absorber material on a Mo-coated glass substrate (preferably soda-lime glass). The properties of the Mo film and the choice of the glass substrate are of primary importance for the final device quality, because of the importance of Na, which diffuses from the glass through the Mo film into the growing absorber material. In the past, some processes used blocking layers such as SiNx, SiO₂, or Cr between the glass substrate and the Mo film to prevent the out-diffusion of Na. Instead, Na-containing precursors like NaF [45], Na₂Se [46], or Na₂S [47] are then deposited prior to absorber growth to provide a controlled, more homogeneous, incorporation of Na into the film. The control of Na incorporation in the film from precursor layers allows the use of other substrates like metal or polymer foils. The most obvious effects of Na incorporation are better film morphology and higher conductivity of the films [48]. Furthermore, the incorporation of Na induces beneficial changes in the defect distribution of the absorber films [49,50].

The explanations for the beneficial impact of Na are manifold, and it is most likely that the incorporation of Na in fact results in a variety of consequences. During film growth, the incorporation of Na leads to the formation of NaSe_x compounds. This slows down the growth of CuInSe₂ and could at same time facilitate the incorporation of Se into the film [51]. Also the widening of the existence range of the α-(CuInSe₂) phase in the phase diagram, discussed above, as well as the reported larger tolerance to the Cu/(In+Ga) ratio of Na-containing thin films, could be explained in this picture. Furthermore, the higher conductivity of Na-containing films could result from the diminished number of compensating V_se donors.

During absorber deposition, a MoSe₂ film forms at the Mo surface [52,53]. MoSe₂ is a layered semiconductor with p-type conduction, a band gap of 1.3 eV and weak van der Waals bonding along the c-axis. The c-axis is found to be in parallel with, and the van der Waals planes thus perpendicular to, the interface [53]. Because of the larger band gap of the MoSe₂ compared with that of standard Cu(In,Ga)Se₂ films, the MoSe₂ layer provides a low-recombinative back surface for the photogenerated minority carriers (electrons) in the Cu(In,Ga)Se₂ absorber and at the same time provides a low-resistance contact for the majority carries (holes).

Photovoltaic-grade Cu(In,Ga)Se₂ films have a slightly In-rich overall composition. The allowed stoichiometry deviations are astonishingly large, yielding a wide process window with respect to composition. Devices with efficiencies above 14% are obtained from absorbers with (In+Ga)/(ln+Ga+Cu) ratios between 52% and 64% if the sample contains Na [48]. Cu-rich Cu(In,Ga)Se₂ shows the segregation of a secondary Cu_ySe phase preferentially at the surface of the absorber film. The metallic nature of this phase does not allow the formation of efficient heterojunctions. Even after removal of the secondary phase from the surface by etching the absorber in KCN, the utility of this material for photovoltaic applications is limited. However, the importance of the Cu-rich composition is given by its role during film growth. Cu-rich films have grain sizes in excess of 1 μm whereas In-rich films have much smaller grains. A model for the film growth under Cu-rich compositions comprises the role of Cu_ySe as a flux agent during the growth process of co-evaporated films [54]. For Cu(In,Ga)Se₂ prepared by selenisation, the role of Cu_ySe is similar [55]. Therefore, growth processes for high quality material have to go through a copper-rich stage but have to end up with an indium-rich overall composition.

3.3.2 Co-Evaporation Processes

The absorber material yielding the highest efficiencies is Cu(In,Ga)Se₂ with a Ga/(Ga + In) ratio of 20–30%, prepared by co-evaporation from elemental sources. Figure 6 sketches a co-evaporation set-up as used for the preparation of laboratory-scale solar cells and mini-modules. The process requires a maximum substrate temperature of about 550 °C for a certain time during film growth, preferably towards the end of growth. One advantage of the evaporation route is that material deposition and film formation are performed during the same processing step. A feedback loop based on a quadrupole mass spectrometer or an atomic absorption spectrometer controls the rates of each source. The composition of the deposited material with regard to the metals corresponds to their evaporation rates, whereas Se is always evaporated in excess. This precise control over the deposition rates allows for a wide range of variations and optimisations with different sub-steps or stages for film deposition and growth. These sequences are defined by the evaporation rates of the different sources and the substrate temperature during the course of deposition.

FIGURE 6 Arrangement for the deposition of Cu(In,Ga)Se₂ films on the laboratory scale by co-evaporation on a heated substrate. The rates of the sources are controlled by mass spectrometry.

Advanced preparation sequences always include a Cu-rich stage during the growth process and end up with an In-rich overall composition in order to combine the large grains of the Cu-rich stage with the otherwise more favourable electronic properties of the In-rich composition. The first example of this kind of procedure is the so called Boeing or bilayer process [5], which starts with the deposition of Cu-rich Cu(In,Ga)Se₂ and ends with an excess In rate. The most successful co-evaporation process is the so-called three-stage process [56]. This process starts with the deposition of In,Ga, and Se at relatively low temperatures, then uses a Cu-rich growth stage by evaporating Cu in excess at elevated temperature, and at the end again deposits only In, Ga, and Se to ensure the overall In-rich composition of the film. The three-stage process currently leads to the Cu(In,Ga)Se₂ solar cells with the highest efficiencies.

3.3.3 Selenisation Processes

The second class of absorber preparation routes is based on the separation of deposition and compound formation into two different processing steps. High efficiencies are obtained from absorber prepared by selenisation of metal precursors in H₂Se [57–59] and by rapid thermal processing of stacked elemental layers in a Se atmosphere [60]. These sequential processes have the advantage that approved large-area deposition techniques like sputtering can be used for the deposition of the materials. The Cu(In,Ga)Se₂ film formation then requires a second step, the selenisation.

The very first large-area modules were prepared by the selenisation of metal precursors in the presence of H₂Se more than ten years ago [6]. Today, a modification of this process yields the first commercially available Cu(In,Ga)Se₂ solar cells manufactured by Shell Solar Industries. This process is schematically illustrated in Figure 7. First, a stacked layer of Cu, In and Ga is sputter-deposited on the Mo-coated glass substrate. Then selenisation takes place under H₂Se. To improve the device performance, a second thermal process under H₂S is added, resulting in an absorber that is Cu(In,Ga)(S,Se)₂ rather than Cu(In,Ga)Se₂.

FIGURE 7 Illustration of the sequential process. First a stack of metal (Cu,In,Ga) layers deposited by sputtering onto a Mo-coated glass. In the second step, this slack is selenised in H₂Se atmosphere and converted into CuIn Se₂.

A variation of the method that avoids the use of the toxic H₂Se during selenisation is the rapid thermal processing of stacked elemental layers [60]. Here the precursor includes a layer of evaporated elemental Se. The stack is then selenised by a rapid thermal process (RTP) in either an inert or a Se atmosphere. The highest efficiencies are obtained if the RTP is performed in an S-containing atmosphere (either pure S or H₂S).

On the laboratory scale, the efficiencies of cells made by these preparation routes are smaller by about 3% (absolute) as compared with the record values. However, on the module level, co-evaporated and sequentially prepared absorbers have about the same efficiency. Sequential processes need two or even three stages for absorber completion. These additional processing steps may counterbalance the advantage of easier element deposition by sputtering.

3.3.4 Other Absorber Deposition Processes

Besides selenisation and co-evaporation, other deposition methods have been studied, either to obtain films with very high quality or to reduce the cost of film deposition on large areas. Methods that are used to grow epitaxial III–V compound films, such as molecular beam epitaxy (MBE) [61] or metal organic chemical vapour deposition (MOCVD) [62] have revealed interesting features for fundamental studies like phase segregation and defect formation, but could not be used to form the absorber material for high-efficiency solar cells.

Attempts to develop so-called low-cost processes include electrodeposition [63–66], screen printing and particle deposition [67,68]. Electrodeposition can be done either in one or two steps. The crucial step is the final film formation in a high-temperature annealing process. The recrystallisation process competes with the decomposition of the material. Therefore, process optimisation is quite difficult. Cells with high efficiencies were obtained by electrodeposition of a Cu-rich CuInSe₂ film and subsequent conditioning by a vacuum evaporation step of In(Se) [69]. Particle deposition by printing of suitable inks and subsequent annealing lead to absorber layers with good quality enabling the fabrication of solar cells with efficiencies over 13% [70].

3.3.5 Postdeposition Air Anneal

Air annealing has been an important process step, crucial for the efficiency especially of the early solar cells based on CuInSe₂. Though often not mentioned explicitly, an oxygenation step is still used for most of the present-day high-efficiency devices. The beneficial effect of oxygen was explained within the defect chemical model of Cahen and Noufi [71]. In this model, the surface defects at grain boundaries are positively charged Se vacancies V_Se. During air annealing, these sites are passivated by O atoms. Because of the decreased charge at the grain boundary, the band bending as well as the recombination probability for photogenerated electrons is reduced. The surface donors and their neutralisation by oxygen are important for the free Cu(In,Ga)Se₂ surface as well as for the formation of the CdS/Cu(In,Ga)Se₂ interface [72,73].

3.4 Heterojunction Formation

3.4.1 The Free Cu(In,Ga)Se₂ Surface

The surface properties of Cu(In,Ga)Se₂ thin films are especially important as this surface becomes the active interface of the completed solar cell. However, the band diagram of the ZnO/CdS/Cu(In,Ga)Se₂ heterojunction, especially the detailed structure close to the CdS/Cu(In,Ga)Se₂ interface, is still under debate (for recent reviews see references [74,75]).

The free surfaces of as-grown Cu(In,Ga)Se₂ films exhibit two prominent features:

1. The valence band-edge energy EV lies below the Fermi level EF by about 1.1 eV for CuInSe₂ films [76]. This energy is larger than the band-gap energy E^bulk_g of the bulk of the absorber material. This finding was taken as an indication for a widening of band gap at the surface of the film. A recent direct measurement of the surface band gap of polycrystalline CuInSe₂ by Morkel et al. [77] proved that the band-gap energy E^surf_g at the surface of the film is about 1.4 cV, i.e., more than 0.3 eV larger than E^bulk_g≈1.04 eV. In Cu(In₁₋_x,Ga_x)Se₂ alloys the distance EF– EV was found to be 0.8 eV (almost independently of the Ga content if x > 0) [78].

2. The surface composition of Cu-poor CuInSe₂, as well as that of Cu(In,Ga)Se₂ films, corresponds to a surface composition of (Ga + In)/(Ga + In + Cu) of about 0.75 for a range of bulk compositions of 0.5 < (Ga + In)/(Ga + In + Cu) < 0.75 [76].

Both observations (i) and (ii) have led to the assumption that a phase segregation of Cu(In,Ga)₃Se₅, the so-called Ordered Defect Compound (ODC), occurs at the surface of the films. From the fact that bulk Cu(In,Ga)₃Se₅ exhibits n-type conductivity [79] it was argued that Cu-poor Cu(In,Ga)Se₂ thin films automatically generate a rectifying, buried junction. However, the existence of a separate phase on top of standard Cu(In,Ga)Se₂ thin films has not yet been confirmed by structural methods such as X-ray diffraction, high resolution transmission electron microscopy or electron diffraction. Furthermore, if the surface phase exhibited the weak n-type conductivity of bulk Cu(In,Ga)₃Se₅, simple charge neutrality estimations [21] show that this would not be sufficient to achieve type inversion.

Based on these arguments, another picture of the surface of Cu(In,Ga)Se₂ thin films and of junction formation has emerged [21,80]. Within the classical Bardeen model [81] of Fermi level pinning by electronic states at semiconductor surfaces, a density of surface states of about 10¹² cm⁻² eV⁻¹ is sufficient to pin the Fermi level at the neutrality level of the free semiconductor surfaces. Positively charged surface donors are expected in the metal terminated (112) surface of CuInSe₂ owing to the dangling bond to the missing Se [71]. Thus, these surface states, rather than a distinct n-type surface phase, determine the type inversion of the surface.

Surface states play also an important role in the completed heterostructure, where they become interface states at the absorber/buffer interlace. The deject layer model [21,80] takes into account a modification of the band structure due to the Cu deficiency of the surface as well as the presence of positively charged surface states due to the missing surface Se. However, the defect layer model considers the surface layer not as n-type bulk material (as does the ODC model) but as a p⁺-layer. Furthermore, the defect layer is viewed not as the origin but rather as the consequence of the natural surface type inversion [21,80]. Surface states are responsible for the surface band bending that leads to the liberation of Cu from its lattice sites and to Cu migration towards the neutral part of the film. The remaining copper vacancies close to the surface result in a high density of acceptor states, i.e., the p⁺-defect layer at the film surface. Recent photoelectron spectroscopy experiments of Klein and Jaegermann [82] suggest that the band bending occurring during junction formation leads to a loss of Cu atoms from the surface of CuInSe₂ and CuGaSe₂, whereas a similar effect was not observed in CuInS₂.

3.4.2 Buffer Layer Deposition

Surface passivation and junction formation is most easily achieved by the CBD deposition of a thin CdS film from a chemical solution containing Cd ions and thiourea [83]. The benefit of the CdS layer is manifold:

• CBD deposition of CdS provides complete coverage of the rough polycrystalline absorber surface at a film thickness of only 10 nm [84].

• The layer provides protection against damage and chemical reactions resulting from the subsequent ZnO deposition process.

• The chemical bath removes the natural oxide from the film surface [83] and thus reestablishes positively charged surface states and, as a consequence, the natural type inversion at the CdS/Cu(In,Ga)Se₂ interface.

• The Cd ions, reacting first with the absorber surface, remove elemental Se, possibly by the formation of CdSe.

• The Cd ions also diffuse to a certain extent into the Cu-poor surface layer of the absorber material [85,86], where they possibly form Cd_Cu donors, thus providing additional positive charges enhancing the type inversion of the buffer/absorber interface.

• Open-circuit voltage limitations imposed by interface recombination can be overcome by a low surface recombination velocity in addition to the type inversion of the absorber surface. Thus, one might conclude that interface states (except those shallow surface donors responsible for the type inversion) are also passivated by the chemical bath.

Due to the favourable properties of CdS as a heterojunction partner and the chemistry of the CBD process it is difficult to find a replacement. Avoiding CdS and the chemical bath step would be advantageous from the production point of view. On one hand, a toxic material such as CdS requires additional safety regulation; on the other hand, the chemical bath deposition does not comply with the vacuum deposition steps of an in-line module fabrication. Therefore, research and development in this area relates to two issues: (i) the search for alternative materials for a chemical deposition, and (ii) the development of ways to deposit the front electrode without an intermediate step in a chemical bath.

Promising materials to replace CdS are In(OH,S) [78], Zn(OH,S) [88,89], ZnS [90,91], and ZnSe [92–94], However, all these materials require additional precautions to be taken for the preparation of the absorber surface or front electrode deposition.

3.4.3 Window Layer Deposition

The most commonly used material for the preparation of the front electrode is ZnO doped with B or Al. The first large-area modules produced by ARCO Solar (later Siemens Solar Industries, now Shell Solar Industries) had a ZnO:B window layer deposited by chemical vapour deposition (CVD). Later production facilities at Boeing and EUROCIS use sputtering processes. Present pilot production lines also favour sputtering. As mentioned previously, an undoped i-ZnO layer with a thickness of about 50–100 nm is needed at the heterojunction in order to achieve optimum performance.

3.5 Module Production and Commercialisation

3.5.1 Monolithic Interconnections

One inherent advantage of thin-film technology for photovoltaics is the possibility of using monolithic integration for series connection of individual cells within a module. In contrast, bulk Si solar cells must be provided with a front metal grid, and each of these front contacts has to be connected to the back contact of the next cell for series connection. The interconnect scheme, shown in Figure 8, has to warrant that the front ZnO layer of one cell is connected to the back Mo contact of the next one. In order to obtain this connection, three different patterning steps are necessary. The first one separates the Mo back contact by a series of periodical scribes and thus defines the width of the cells, which is of the order of 0.5–1 cm. For Mo patterning, a laser is normally used. The second patterning step is performed after absorber and buffer deposition, and the final one after window deposition (cf. Figure 8). Scribing of the semiconductor layer is done by mechanical scribing or laser scribing. The total width of the interconnect depends not only on the scribing tools, but also on the reproducibility of the scribing lines along the entire module. The length of the cells and, accordingly, that of the scribes can be more than 1 m. The typical interconnect width is of the order of 300 μm. Thus, about 3–5% of the cell area must be sacrificed to the interconnects.

FIGURE 8 Deposition and patterning sequence to obtain an integrated interconnect scheme for Cu(In,Ga)Se₂ thin-film modules.

3.5.2 Module Fabrication

The technologies for absorber, buffer, and window deposition used for module production are the same as those discussed above for the production of small laboratory cells. However, the challenge of producing modules is to transform the laboratory-scale technologies to much larger areas. The selenisation process uses as much off-the-shelf equipment and processing as possible (e.g., sputtering of the metal precursors) for fabricating Cu(In,Ga)Se₂ films. For co-evaporation on large areas, the Centre for Solar Energy and Hydrogen Research in Stuttgart (ZSW) has designed its own equipment, schematically shown in Figure 9, for an in-line co-evaporation process. Line-shaped evaporation sources allow continuous deposition of large-area, high-quality Cu(In,Ga)Se₂ films. The relatively high substrate temperatures that are necessary for high-quality material impose problems in handling very large area glass sheets. Future process optimisation therefore implies reduction of the substrate temperature.

FIGURE 9 Sketch of an in-line deposition system for co-evaporation of Cu(In,Ga)Se₂ absorber films from line-sources.

A bottleneck for the production is the deposition of the buffer layer in a chemical bath. On one hand, it is not straightforward to integrate this process in a line consisting mainly of dry physical vapour deposition processes; on the other, it would be favourable for environmental reasons to replace the currently used CdS layer by a Cd-free alternative.

The ZnO transparent front electrode is put down either by CVD or sputtering. Each method has its specific advantages with respect to process tolerance, throughput, cost and film properties. The widths of the cells within the module—and therefore the relative losses from the patterning—mainly depend on the sheet resistance of the ZnO.

Module encapsulation is an important issue because module stability depends on proper protection against humidity. Low-iron cover glasses provide good protection against ambient influences. Hermetic sealing of the edges is mandatory to obtain stable modules (see the following discussion).

3.5.3 Upscaling Achievements

Cu(In,Ga)Se₂ has the best potential to reach more than 15% module efficiency in the near future. Mini-modules ranging in area from 20 cm² to 90 cm² that use the process sequence anticipated for larger area commercial modules have already reached efficiencies around 14–15%. Recently, Siemens Solar Industries fabricated a 1 ft × 4 ft power module (−44 watts) with an independently verified efficiency of 12.1%. Using a totally different approach to the deposition of the absorber layer, the ZSW fabricated a 30-cm × 30-cm module with a verified efficiency of 12.7%. Based on the same co-evaporation process, Würth Solar GmbH, Stuttgart, reported an efficiency of 12.5% for a module of aperture area 5932 cm² [95]. More results of the different processes are summarised in Table 2. Note that, because of the promising results from the laboratory scale and the first approaches of up-scaling, several companies other than those mentioned in Table 2 now plan commercial production.

TABLE 2 Comparison of Efficiencies η and Areas A of Laboratory Cells, Mini-Modules, and Commercial-Size Modules Achieved with Cu(In,Ga)Se₂ Thin Films Based on the Co-Evaporation and the Selenisation Process. NREL Denotes the National Renewable Energy Laboratories (USA), ZSW is the Center for Solar Energy and Hydrogen Research (Germany), EPV is Energy Photovoltaics (USA), ASC is the Angstrom Solar Centre (Sweden)

^aFlexible cells.

A further challenge is to develop CIGS cells on flexible substrates and hence to extend their area of applications. There are ongoing efforts to produce cells on various kinds of substrates like stainless steel, polymide and at the same time to retain the performance achieved with devices on soda-lime glass [96]. For space applications it is important to reduce the weight by depositing the cells on lightweight foil substrates. Highest small-area efficiencies on polymide films formed by spin coating on a glass substrate reach 12.8% [97]. Roll to roll coating on metal foils [98] and polymer films [99,100] has already reached the stage of pilot production.

3.5.4 Stability

The long-term stability is a critical issue of any solar cell technology because the module lifetime contributes as much to the ratio between produced energy and invested cost as does the initial efficiency. Cu(In,Ga)Se₂ modules fabricated by Shell Solar Industries more than 10 years ago show until today very good stability during outdoor operation [101,102]. However, intense accelerated lifetime testing is made for the now commercially available Cu(In,Ga)Se₂ modules. Careful sealing and encapsulation appears mandatory, especially because of the sensitivity of Cu(In,Ga)Se₂ to humidity. For non-encapsulated modules, corrosion of the molybdenum contact and the degradation of zinc oxide were found to be the dominating degradation mechanisms [103] during the so-called damp heat test (1000 hours in hot (85 °C) and humid (85% humidity) atmosphere). Investigations of non-encapsulated cells [104–106] unveiled further a humidity-induced degradation of the Cu(In,Ga)Se₂ absorber material. Despite of the sensibility of Cu(In,Ga)Se₂ with respect to humidity, well encapsulated modules pass the damp heat test [107].

Recent work of Guillemoles et al. [108,109] investigates the chemical and electronic stability of Cu(In,Ga)Se₂-based solar cells and possible fundamental instabilities of the material system, namely, interface reactions, defect metastability, and constituent element (Cu) mobility. Guillemoles et al. conclude that all reasonably anticipated detrimental interface reactions at the Mo/Cu(In,Ga)Se₂, the Cu(In,Ga)Se₂/CdS. or the CdS/ZnO interface are either thermodynamically or kinetically limited, Furthermore, Cu mobility does not contradict long-term stability [108,109].

3.5.5 Radiation Hardness and Space Applications

One important prospective application for Cu(In,Ga)Se₂ cells is in space, where the main power source is photovoltaics. Satellites in low-earth orbits for communication systems require solar cells with high end-of-life efficiencies, despite the high flux of high-energy electrons and protons in that ambient. The radiation hardness of Cu(In,Ga)Se₂ has been recognized as early as 1984/85 [110,111], but only recently have systematic investigations on the radiation response of Cu(In,Ga)Se₂ solar cells been undergone using high enough electron and proton fluences to allow quantitative conclusions regarding the defect generation rates [112]. The radiation resistance of Cu(In,Ga)Se₂ against high-energy (0.3–3 MeV) electrons turns out to be far better than that of any other photovoltaic material [112]. The radiation hardness of Cu(In,Ga)Se₂ against high-energy (0.4–10 MeV) protons is also high, though the difference from other materials is not as high as in case of electron irradiation. Walters et al. [113] have analysed Cu(In,Ga)Se₂ in the frame of the so-called damage dose model and found that the damage coefficients for Cu(In,Ga)Se₂ are comparable to those of InP and considerably smaller than those of Si and GaAs.

The high mobility of Cu in the Cu(In,Ga)Se₂ lattice was proposed to be one important ingredient for a defect healing mechanism that could explain the high radiation resistance of Cu(In,Ga)Se₂ [108,114]. Recent thermal annealing experiments of electron-irradiated [115] and proton-irradiated [116] Cu(In,Ga)Se₂ solar cells unveiled a thermally activated healing process with an activation energy of around 1 cV leading to a complete recovery of the device performance. Illumination of the solar cell enhances this annealing process further [116,117].

The challenge for developing CIGS space cells is to reduce the weight by depositing the cells on foil substrates, and at the same time to retain the performance achieved with devices on soda-lime glass. Recently, Tuttle et al. [118] reported a Cu(In,Ga)Se₂ solar cell on lightweight metal foil with a power conversion efficiency of 15.2% (under AM0 illumination) and a specific power of 1235 W/kg. Other approaches to flexible and lightweight Cu(In,Ga)Se₂ solar cells embrace Cu(In,Ga)Se₂ deposited on plastic foil such as polyimide [119] or the use of a lift-off technique to remove the absorber from the glass substrate after device fabrication [120].

4 Device Physics

4.1 The Band Diagram

The band diagram of the ZnO/CdS/Cu(In,Ga)Se₂ heterostructure in Figure 10 shows the conduction and valence band energies E_c and E_v of the Cu(In,Ga)Se₂ absorber, the CdS buffer layer and the ZnO window. Note that the latter consists of a highly Al-doped (ZnO:Al) and an undoped (i-ZnO) layer. For the moment, we neglect the polycrystalline nature of the semiconductor materials, which in principle requires a two- or three-dimensional band diagram. Even in the one-dimensional model, important details of the band diagram are still not perfectly clear. The diagram in Figure 10 concentrates on the heterojunction and does not show the contact between the Mo and Cu(In,Ga)Se₂ at the back side of the absorber.

FIGURE 10 Band diagram of the ZnO/CdS/Cu(In,Ga)Se₂ heterojunction under bias voltage showing the conduction and valence band-edge energies ΔE_c and E_v. The quantities ΔE_c^wb/ba denote the conduction band offsets at the window/buffer and buffer/absorber interfaces, respectively. An internal valence band offset ΔE_v^int exists between the bulk Cu(In,Ga)Se₂ and a surface defect layer (SDL) on top of the Cu(In,Ga)Se₂ absorber film. The quantity ΔE_Fn denotes the energy distance between the electron Fermi level E_Fn and the conduction band at the CdS buffer/Cu(In,Ga)Se₂ absorber interface, and ϕ_n denotes the neutrality level of interface states at this heterointerface.

An important feature in Figure 10 is the 10–30-nm-thick surface defect layer (SDL) on top of the Cu(In,Ga)Se₂ absorber, already discussed in Section 3.4. The physical nature of this SDL is still under debate. However, the fact that the band gap at the surface of Cu(In,Ga)Se₂ thin films (as long as they are prepared with a slightly Cu-poor final composition) exceeds the band-gap energy of the bulk material [76,77], has important implications for the contribution of interface recombination to the overall performance of the solar cell. A simplified approach to the mathematical description of the ZnO/CdS/Cu(In,Ga)Se₂ heterojunction including the consequence of the surface-band-gap widening is given in reference [121].

The most important quantities to be considered in the band diagram are the band discontinuities between the different heterojunction partners. Band discontinuities in terms of valence band offsets ΔI between semiconductor a and b are usually determined by photoelectron spectroscopy (for a discussion with respect to Cu-chalcopyrite surfaces and interfaces, see [74]). The valence band offset between a (011)-oriented Cu(In,Ga)Se₂ single crystal and CdS deposited by PVD at room temperature is determined as [122,123] ΔE ^ab_v = −0.8 (±0.2) eV (and, therefore, ΔE^ab_c = E^bCdS_g – E^CIS_g + ΔE^ab_v≈0.55 eV, with the band gaps E^CdS_g≈2.4 eV and E^CdS_g≈1.05 eV of CdS and CuInSe₂, respectively). Several authors have investigated the band discontinuity between polycrystalline Cu(In,Ga)Se₂ films and CdS, and found values between − 0.6 and − 1.3 eV with a clear centre of mass around − 0.9 eV, corresponding to a conduction band offset of 0.45 eV [74]. Wei and Zunger [124] calculated a theoretical value of ΔE^ab_v = −1.03 eV, which would lead to ΔE^ab_C≈0.3 eV. Recently, Morkel et al. [77] found a valence band offset between the surface of polycrystalline, Cu-poor prepared CuInSe₂ and chemical bath deposited CdS of ΔE^ab_c = −0.8 eV. By combining their photoelectron spectroscopy results with measured surface band-gap energies of CdS and CuInSe₂ from inverse photoemission spectroscopy, the authors of reference [77] conclude that the conduction band offset ΔE^ab_c is actually zero. This is because the deposited CdS has E_g = 2.2 eV due to S/Se intermixing and the CuInSe₂ film exhibits a surface band gap of 1.4 eV; thus, ΔE^ab_c = Eg[Cd(Se, S)] – E^surf_gs [CuInSe₂] + ΔE^ab_V ≈ 0.

The band alignment of polycrystalline CuInSe₂ and Cu(In,Ga)Se₂ alloys was examined by Schmid et al. [76,78] who found that the valence band offset is almost independent of the Ga content. In turn, the increase of the absorber band gap leads to a change of ΔE^ab_C from positive to negative values. The conduction band offset between the CdS buffer and the ZnO window layer was determined by Ruckh et al. to be 0.4 eV [125].

4.2 Short-Circuit Current

The short-circuit current density J_se that can be obtained from the standard 100 mW cm⁻² solar spectrum (AM1.5) is determined, on one hand, by optical losses, that is. by the fact that photons from a part of the spectrum are either not absorbed in the solar cell or are absorbed without generation of electron–hole pairs. On the other hand, not all photogenerated electron–hole pairs contribute to J_sc because they recombine before they are collected. We denote these latter limitations as recombination losses. Figure 11 illustrates how much from an incoming photon flux from the terrestrial solar spectrum contributes to the final J_sc of a highly efficient Cu(In,Ga)Se₂ solar cell [126] and where the remainder gets lost. The incoming light, i.e., that part of the solar spectrum with photon energy hv larger than the band-gap energy E_g = 1.155 eV of the specific absorber would correspond to a (maximum possible) J_sc of 41.7 mA cm^–2. By reflection at the surface and at the three interfaces between the MgF₂ anti-reflective coating, the ZnO window, the CdS buffer, and the Cu(In,Ga)Se₂ absorber layer we loose already 0.6 mA cm⁻². A further reflection loss of 0.1 mA cm⁻² is due to those low-energy photons that are reflected at the metallic back and leave the solar cell after having traversed the absorber twice. Due to the high absorption coefficient of Cu(In,Ga)Se₂ for hv > E_g, (this portion is very small. More important are parasitic absorption losses by free-carrier absorption in the highly doped part of the ZnO window layer (0.9 mA cm⁻²) and at the absorber/Mo interface (0.5 mA cm⁻²). Thus, the sum of all optical losses amounts to 2.1 mA cm⁻². Note that in solar cells and modules there are additional optical losses due to the grid shadowing or interconnect areas, respectively. These losses are not discussed here.

FIGURE 11 Optical and electronic losses of the short circuit current density J_sc of a high-efficiency ZnO/CdS/Cu(In,Ga)Se₂ heterojunction solar cell. The incident current density of 41.7 mA/cm² corresponds to the range of the AM 1.5 solar spectrum that has a photon energy larger than the hand gap energy E_g = 1.155 eV of the Cu(In,Ga)Se₂ absorber. Optical losses consist of reflection losses at the ambient/window, at the window/buffer, the buffer/absorber, and at the absorber/back contact interface as well as of parasitic absorption in the ZnO window layer (free carrier absorption) and at the Mo back contact. Electronic losses are recombination losses in the window, buffer, and in the absorber layer. The finally measured J_sc of 34.6 mA/cm² of the cell stems almost exclusively from the Cu(In,Ga)Se₂ absorber and only to a small extend from the CdS buffer layer.

Next, we have to consider that electron hole pairs photogenerated in the ZnO window layer are not separated. Therefore, this loss affecting photon energies hv > E_g (ZnO) = 3.2 eV contributes to the recombination losses. As shown in Figure 11, this loss of high-energy photons costs about 0.7 mA cm⁻². Another portion of the solar light is absorbed in the CdS buffer layer (E_g (CdS)≈2.4 eV). However, a part of the photons in the energy range 3.2 eV > hv > 2.4 eV contributes to J_sc because the thin CdS layer does not absorb all those photons and a part of the electron–hole pairs created in the buffer layer still contributes to the photocurrent [127]. In the present example 1.3 mA cm⁻², get lost by recombination and 0.8 mA cm⁻² are collected. However, the major part of J_sc (33.8 mA cm⁻²) stems from electron–hole pairs photogenerated in the absorber. Finally, the collection losses in the absorber amount to 3.2 mA cm⁻².

The preceding analysis shows that, accepting the restrictions that are given by the window and the buffer layer, this type of solar cell makes extremely good use of the solar spectrum. There is however still some scope for improving J_sc by optimising carrier collection in the absorber and/or by replacing the CdS buffer layer by an alternative material with a higher E_g.

4.3 Open-Circuit Voltage

At open circuit no current flows across the device, and all photogenerated charge carriers have to recombine within the solar cell. The possible recombination paths for the photogenerated charge carriers in the Cu(In,Ga)Se₂ absorber are indicated in the band diagram of Figure 12. Here we have considered recombination at the back surface of the absorber (A´) and in the neutral bulk (A), recombination in the space-charge region (B), and recombination at the buffer/absorber interface (C). Note that due to the presence of high electrical fields in the junction region, the latter two mechanisms may be enhanced by tunnelling.

FIGURE 12 Recombination paths in a ZnO/CdS/(low-gap) Cu(In,Ga)Se2 junction at open circuit. The paths A represent recombination in the neutral volume, A recombination at the back contact, B recombination in the space-charge region, and C recombination at the interface between the Cu(In,Ga)Se₂ absorber and the CdS buffer layer. Back contact recombination is reduced by the conduction band offset ΔE_c^back between the Cu(In,Ga)Se₂ absorber and the MoSe2 layer that forms during absorber preparation on top of the metallic Mo back contact. Interface recombination (C) is reduced by the internal valence band offset ΔE_v^int between the bulk of the Cu(In,Ga)Se2 absorber and the Cu-poor surface layer. The quantity ϕ∗_bp denotes the energy barrier at the CdS/absorber interface and E_T indicates the the energy of a recombination centre in the bulk of the Cu(In,Ga)Se₂.

The basic equations for the recombination processes (A–C) can be found, for example, in [128]. Notably, all recombination current densities J_R can be written in the form of a diode law

(1)

where V is the applied voltage, n_id the diode ideality factor, and k_BT/q the thermal voltage. The saturation current density J₀ in general is a thermally activated quantity and may be written in the form

(2)

where E_a is the activation energy and the prefactor J₀₀ is only weakly temperature dependent.

The activation energy E_a for recombination at the back surface, in the neutral zone, and in the space charge recombination is the absorber band-gap energy E_g, whereas in case of interface recombination E_a equals the barrier ϕ_bp that hinders the holes from the absorber to come to the buffer/absorber interface (cf. Figure 12). In the simplest cases, the diode ideality factor n_id is unity for back surface and neutral zone recombination as well as for recombination at the buffer/absorber interface, whereas for space charge recombination n_id = 2. Equalising the short-circuit current density J_sc and the recombination current density J_R in Equation (1) for the open-circuit-voltage situation (i.e., V = V_oc in Equation (1)) we obtain with the help of Equation (2)

(3)

Note that Equation (3) yields the open circuit voltage in a situation where there is single mechanism clearly dominating the recombination in the specific device. Note that Equation (3) is often used for the analysis of the dominant recombination path. After measuring V_oc. at various temperatures T, the extrapolation of the experimental V_oc(T) curve to T = 0 yields the activation energy of the dominant recombination process, e.g., V_oc(0) = Eg/q in case of bulk recombination or V_oc(0) = ϕ_bp/q in case of interface recombination.

We emphasise that, in practice, measured ideality factors considerably deviate from that textbook scheme and require more refined theories (see, e.g., [14] and references therein). However, for a basic understanding of the recombination losses in thin-film solar cells, a restriction to those textbook examples is sufficient.

The band diagram in Figure 12 contains two important features that appear important for the minimisation of recombination losses in Cu(In,Ga)Se₂ solar cells. The first one is the presence of a considerably widened band gap at the surface of the Cu(In,Ga)Se₂ absorber film. This surface-band-gap widening implies a valence band offset ΔE^int_v at the internal interface between the absorber and the surface defect layer (discussed in Section 4.1). This internal offset directly increases the recombination barrier ϕ_bp to an effective value

(4)

Substituting Equation (4) in Equation (3) we obtain

(5)

It is seen that the internal band offset directly adds to the open-circuit voltage that is achievable in situations where only interface recombination is present. However, the open circuit voltage of most devices that have this surface-band-gap widening (those having an absorber with a final Cu-poor composition) is then limited by bulk recombination.

The second important feature in Figure 12 is the conduction band offset ΔE^back_c between the Cu(In,Ga)Se₂ absorber and the thin MoSe₂ film. The back surface recombination velocity at the metallic Mo back contact is reduced from the value S_b (without the MoSe₂) to a value S^∗_b = S_b exp(−ΔE_c/k_BT).

4.4 Fill Factor

The fill factor FF of a solar cell can be expressed in a simple way as long as the solar cell is well described by a diode law. Green [129] gives the following phenomenological expression for the fill factor

(6)

It is seen that, through the dimensionless quantity v_oc = qV_oc/n_idk_BT, the fill factor depends on the temperature T as well as on the ideality factor n_id of the diode (see also Chapter Ia-1). The fill factor FF₀ results solely from the diode law form of Equation (1). In addition, effects from series resistance R_s and shunt resistance R_p add to the fill factor losses. A good approximation is then given by

(7)

with the normalised series and parallel resistances given by r_s = R_sJ_sc/V_oc and r_p = R_pJ_sc/V_oc, respectively. The description of Cu(In,Ga)Se₂ solar cells in terms of Equations (6) and (7) works reasonably well, e.g., the world record cell [1] has a fill factor of 78.6% and the value calculated from Equations (6) and (7) is 78.0% (V_oc = 678 mV, J_sc = 35.2 mA cm⁻², R_s = 0.2 Ω cm², R_p = 10⁴ Ω cm², n_id = 1.5).

Factors that can further affect the fill factor are (i) the voltage bias dependence of current collection [130], leading to a dependence of J_sc(V ) on voltage V in Equation (1), (ii) recombination properties that are spatially inhomogeneous [131], or (iii) unfavourable band offset conditions at the heterointerface [132].

4.5 Electronic Metastabilities

The long time increase (measured in hours and days) of the open-circuit voltage V_oc of Cu(In,Ga)Se₂ based solar cells during illumination is commonly observed phenomenon [133,134]. In some cases it is not only V_oc but also the fill factor that improves during such a light soaking procedure. Consequently, light soaking treatments are systematically used to re-establish the cell efficiency after thermal treatments [107,135].

For the present day ZnO/CdS/Cu(In,Ga)Se₂ heterojunctions it appears established that there exist at least three types of metastablities with completely different fingerprints [136,137].

Type I: A continuous increase of the open circuit voltage during illumination and the simultaneous increase of the junction capacitance. Both phenomena are satisfactorily explained as a consequence of persistent photoconductivity in the Cu(In,Ga)Se₂ absorber material [138,139], i.e., the persistent capture of photogenerated electrons in traps that exhibit large lattice relaxations like the well-known DX centre in (Al,Ca)As [140]. This type of metastability affects exclusively the open circuit voltage and can vary from few mV up to 50 mV, especially if the sample has been stored in the dark at elevated temperatures.

Type II: A decrease of the fill factor after the cell has been exposed to reverse voltage bias. In extreme cases, this type of metastability leads to a hysteresis in the IV characteristics, e.g., the fill factor of an illuminated IV curve becomes dependent on whether the characteristic has been measured from negative voltages towards positive ones or vice versa. The application of reverse bias also leads to a metastable increase of the junction capacitance and to significant changes of space charge profiles as determined from capacitance vs. voltage measurements [136,137]. The type II metastability is especially important for devices with non-standard buffer/window combinations (e.g., Cd-free buffer layers) [141,142].

Type III: An increase of the fill factor upon illumination with light that is absorbed in the buffer layer or in the extreme surface region of the Cu(In,Ga)Se₂ absorber, i.e., the blue part of the solar spectrum. This type of instability counterbalances, to a certain extend, the consequences of reverse bias; i.e., it restores the value of the fill factor after it has been degraded by application of reverse bias.

Our overall understanding of metastabilities, especially of type II and III, is still incomplete. Fortunately, all metastabilities observed in Cu(In,Ga)Se₂ so far, tend to improve the photovoltaic properties as soon as the device is brought under operating conditions.

5 Wide-Gap Chalcopyrites

5.1 Basics

The alloy system Cu(In,Ga,Al)(S,Se)₂ provides the possibility of building alloys with a wide range of band-gap energies E_g between 1.04 eV for CuInSe₂ up to 3.45 eV for CuAlS₂ (cf. Figure 2). The highest efficiency within the chalcopyrite system is achieved with the relatively low-band-gap energy E_g of 1.12 eV, and attempts to maintain this high efficiency level at E_g > 1.3 eV have so far failed (for a recent review, see [143]). Practical approaches to wide-gap Cuchalcopyrites comprise (i) alloying of CuInSe₂ with Ga up to pure CuGaSe₂ with E_g = 1.68 eV; (ii) Cu(In,Al)Se₂ alloys with solar cells realised up to an Al/(Al + In) ratio of 0.6 and E_g≈1.8 eV [144]; (iii) CuIn(Se,S)₂ [145], and Cu(In,Ga)(Se,S)₂ [146] alloys with a S/(Se + S) ratio up to 0.5; and (iv) CuInS₂ [147] and Cu(In,Ga)S₂ [148] alloys. Note that the approaches (i)–(iii) are realised with a final film composition that is slightly Cu-poor, whereas approach (iv) uses films that are Cu-rich. In the latter case, CuS segregates preferably at the film surface. This secondary phase has to be removed prior to heterojunction formation.

The advantage of higher voltages of the individual cells by increasing the band gap of the absorber material is important for thin-film modules. An ideal range for the band gap energy would be between 1.4 eV and 1.6 eV because the increased open circuit voltage and the reduced short circuit current density would reduce the number of necessary scribes used for monolithic integration of the cells into a module. Also, the thickness of front and back electrodes can be reduced because of the reduced current density.

Table 3 compares the solar cell output parameters of the best chalcopyrite-based solar cells. This compilation clearly shows the superiority of Cu(In,Ga)Se₂ with a relatively low Ga content which leads to the actual world champion device. The fact that the best CuInSe₂ device has an efficiency of 3% below that of the best Cu(In,Ga)Se₂ device is due not only to the less favourable band-gap energy but also to the lack of the beneficial effect of small amounts of Ga on the growth and on the electronic quality of the thin film, as discussed previously.

TABLE 3 Absorber Band-Gap Energy E_g, Efficiency η, Open-Circuit Voltage V_oc, Short-Circuit Current Density J_sc, Fill Factor FF, and Area A of the Best Cu(In,Ga)Se₂, CuInSe₂, CuGaSe₂, Cu(I-n,Al)Se₂, CuInS₂, Cu(In,Ga)S₂, and Cu(In,Ga)(S,Se)₂ Solar Cells

^aConfirmed total area values.

^bEffective area values (not confirmed).

The difficulty of obtaining wide-gap devices with high efficiencies is also illustrated by plotting the absorber band-gap energies of a series of chalcopyrite alloys vs. the attained open-circuit voltages. Figure 13 shows that below E_g = 1.3 eV, the V_OC data follow a straight line, indicating the proportional gain of V_oc = E_g/q – 0.5 V, whereas at E_g > 1.3 eV the gain is much more moderate. At the high band-gap end of the scale the differences of the band-gap energies and the open-circuit voltages of CuInS₂ and CuGaSc₂ amount to 840 mV and 820 mV, respectively, whereas E_g – qV_oc is only 434 eV in the record Cu(In,Ga)Se₂ device.

FIGURE 13 Open-circuit voltages of different Cu-chalcopyrite bused solar cells with various band-gap energies of the absorber layers. Full symbols correspond to Cu(In,Ga)Se₂ alloys prepared by a simple single layer process (squares), a bilayer process (triangles down), and the three-stage process (triangles up). Cu(In,Ga)Se₂ cells derived from an in-line process as sketched in Fig. 9 are denoted by diamonds. Open triangles relate to Cu(In,Ga)S₂, open circles to Cu(In,Ga)(S,Se)₂, and the crossed triangles to Cu(In, Al)Se₂ cells.

One reason for these large differences E_g – qV_oc in wide-gap devices is the less favourable band offset constellation at the absorber/CdS-buffer interface. Figure 14 shows the band diagram of a wide-gap Cu chalcopyrite-based heterojunction with (Figure 14(a)) and without (Figure 14(b)) the surface defect layer. As the increase of band gap in going from CuInSe₂ to CuGaSe₂ takes place almost exclusively by an increase of the conduction band energy E_c, the positive or zero band offset ΔE^ab_c between the absorber and the buffer of a low-gap device (cf. Figure 12) turns into a negative one in Figure 14. This effect should be weaker for CuInS₂ and CuAlSe₂ as the increase of E_g in these cases is due to an upwards shift of E_c and a downwards shift of E_v. However, any increase of E_c implies that the barrier ϕ_bp that hinders the holes from the absorber from recombination at the heterointerface does not increase proportionally with the increase of the band-gap energy. Thus, in wide-gap absorbers, the importance of interface recombination (determined by the barrier (ϕ_bp) grows considerably relative to that of bulk recombination (determined by E_g of the absorber) [149]. Using the same arguments with respect to the MoSe₂/absorber, the back-surface field provided by this type of heterojunction back contact, as shown in Figure 12, vanishes when the conduction-band energy of the absorber is increased.

FIGURE 14 Energy band diagram of a ZnO/CdS/(wide-gap) Cu(In,Ga)(Se,S)₂ heterojunction. The band diagram (a) that includes the surface defect layer (SDL) of a Cu-poor prepared film shows that the interface recombination barrier ϕ∗_bp = ϕ_bp + ΔE^int_v is larger than the barrier ϕ_bp in the device that was prepared Cu-rich (b). The difference is the internal valence band offset ΔE^int_v between the SDL and the bulk of the absorber. The larger value of ϕ∗_bp reduces interface recombination.

Up to now, all arguments for explaining the relatively low performance of wide-gap chalcopyrite alloys deal with the changes in the band diagram. However, recent work focuses on the differences in the electronic quality of the absorber materials. Here, it was found that the concentration of recombination active defects increases when increasing the Ga content in Cu(In,Ga)Se₂ above a Ga/(Ga + In) ratio of about 0.3 [32] and the S content in Cu(In,Ga)(Se,S)₂ alloys over a S/(Se + S) ratio of 0.3 [34]. Moreover, a recent systematic investigation of the dominant recombination mechanism in large series of Cu(In,Ga)(Se,S)₂ based solar cells with different compositions [150] suggests that bulk recombination prevails in all devices that were prepared with a Cu-poor final film composition. Only devices that had a Cu-rich composition (before removing Curich secondary phases) showed signatures of interface recombination. The absence of interface recombination in Cu-poor devices could be a result of the presence of the Cu-poor surface defect layer in these devices. Comparison of Figures 14(a) and (b) illustrates that interface recombination is much more likely to dominate those devices that lack this feature.

5.2 CuGaSe₂

CuGaSe₂ has a band gap of 1.68 eV and therefore would represent an ideal partner for CuInSe₂ in an all-chalcopyrite tandem structure. However, a reasonable efficiency for the top cell of any tandem structure is about 15%, far higher than has been reached by the present polycrystalline CuGaSe₂ technology. The record efficiency of CuGaSe₂ thin-film solar cells is only 8.3% (9.3% active area) [151] despite of the fact that the electronic properties of CuGaSe₂ are not so far from those of CuInSe₂. However, in detail, all the differences quantitatively point in a less favourable direction. In general, the net doping density N_A in CuGaSe₂ appears too high [152]. Together with the charge of deeper defects, the high doping density leads to an electrical field in the space-charge region, which enhances recombination by tunnelling [153]. The high defect density in CuGaSe₂ thin films absorbers additionally leads to a low diffusion length and, in consequence, to a dependence of the collected short circuit current density J_sc on the bias voltage V. Because of the decreasing width of the space charge region, J_sc(V) decreases with increasing V affecting significantly the fill factor of the solar cell [130]. Note that, on top of the limitation by the unfavourable bulk properties, interface recombination also plays a certain role in CuGaSe₂ [154] because of the unfavourable band diagram shown in Figure 14. Therefore, substantial improvement of the performance requires simultaneous optimisation of bulk and interface properties. Notably, CuGaSe₂ is the only Cu-chalcopyrite material where the record efficiency of cells based on bulk crystals (with η = 9.4%, total area) [115] exceeds that of thin-film devices.

5.3 Cu(In,Al)Se₂

As can be seen in Figure 2, the bandgap change within the Cu(In,Al)Se₂ alloy system is significant even if a small amount of Al is added to CuInSe₂ [144]. This fact allows to grow graded structures with only small changes in the lattice constant. Al–Se compounds tend to react with water vapour to form oxides and H₂Se. Furthermore, there is a tendency to phase segregations [144]. However, cells with very good performance have been achieved by a co-evaporation process of absorbers in a band gap range between 1.09 and 1.57 eV, corresponding to Al/(In + Al) ratio x between 0.09 and 0.59 [144,156]. The highest confirmed efficiency of a Cu(In,Al)Se₂ solar cell is 16.9% [157] (see also Table 3). This device has about the same band-gap energy as the record Cu(In,Ga)Se₂ device [1].

5.4 CuInS₂ and Cu(In,Ga)S₂

The major difference between CuInS₂ and Cu(In,Ga)Se₂ is that the former cannot be prepared with an overall Cu-poor composition. Cu-poor CuInS₂ displays an extremely low conductivity, making it almost unusable as a photovoltaic absorber material [145]. Even at small deviations from stoichiometry on the In-rich side, segregation of the spinel phase is observed [158]. Instead, the material of choice is Cu-rich CuInS₂. As in the case of CuInSe₂, a Cu-rich preparation route implies the removal of the unavoidable secondary Cu–S binary phase by etching the absorber in KCN solution [147]. Such an etch may involve some damage of the absorber surface as well as the introduction of shunt paths between the front and the back electrode. However, as shown in Table 3, the best CuInS₂ device [159] has an efficiency above 11%. This record efficiency for CuInS₂ is achieved by a sulphurisation process rather than by co-evaporation.

As we have discussed previously, interface recombination dominates the open circuit voltage V_oc of Cu(In,Ga)(Se,S)₂ devices that are prepared with a Cu-rich absorber composition (prior to the KCN etch), like the CuInS₂ and Cu(In,Ga)S₂ devices discussed here. It was found recently, that alloying CuInSe₂ with moderate amounts of Ga enhances the open circuit voltage V_oc[160]. This increase of V_or can counterbalance the loss of short circuit current density J_sc resulting from the increased band gap. Apparently, addition of Ga to CuInS₂ reduces interface recombination by increasing the interfacial barrier ϕ_bp(cf. Figure 14(b)) [150]. However, for Cu(In,Ga)S₂ devices prepared by the sulphurisation route, the benefit of Ga alloying is limited, because, during preparation, most of the Ga added to the precursor ends up confined to the rear part of the absorber layer and, therefore, remains ineffective at the absorber surface [160]. In contrast, when preparing Cu(In,Ga)S₂ with co-evaporation, Ga is homogeneously distributed through the depth of the absorber and the Ga content at the film surface is well controlled. Recent work of Kaigawa et al. [148] represents a major progress in wide-gap Cu-chalcopyrites with the preparation of Cu(In,Ga)S₂ solar cells with a confirmed efficiency of 12.3% at a band gap energy E_g = 1.53 eV. In the same publication [148] an efficiency of 10.1% is reported for a device with E_g = 1.65 eV. The open circuit voltage V_oc of this device is 831 mV, i.e., only slightly lower than V_oc of the best CuGaSe₂ cell having an efficiency of 8.3%.

5.5 Cu(In,Ga)(Se,S)₂

One possible way of overcoming the disadvantages of the ternary wide-gap materials CuInS₂ and CuGaSe₂ is to use the full pentenary alloy system Cu(In_1–_xGa_x)(Se_1–_yS_y)₂[146]. Among the materials listed in Table 3, the pentenary system is the only one with an open-circuit voltage larger than 750 mV and an efficiency above 13%, outperforming CuInS₂ in both these respects. The advantage of Cu(In,Ga)(Se,S)₂ could be due to the mutual compensation of the drawbacks of CuGaSe₂ (too high charge density) and that of (Cu-poor) CuInS₂ (too low conductivity, if prepared with a Cu-poor film composition).

5.6 Graded-Gap Devices

An interesting property of the Cu(In,Ga,Al)(S.Se)₂ alloy system is the possibility of designing graded-gap structures in order to optimise the electronic properties of the final device [161–164]. Such band-gap gradings are achieved during co-evaporation by the control of the elemental sources, but selenisation/sulphurisation processes also lead to beneficial compositional gradings. The art of designing optimum band-gap gradings is to push back charge carriers from critical regions, i.e., regions with high recombination probability within the device. Such critical regions are (i) the interface between the back contact and the absorber layer and (ii) the heterojunction interface between the absorber and the buffer material. Figure 15 shows a band diagram of a graded structure that fulfils the requirements for minimising recombination losses.

1. To keep the back contact region clear from electrons, one can use a Ga/In grading. The increase of the Ga/(Ga + In) ratio causes a movement of the conduction band minimum upward with respect to its position in a non-graded CuInSe₂ device. Thus, back surface grading leads to a gradual increase of the conduction-band energy, as illustrated in Figure 15, and therefore drives photogenerated electrons away from the metallic back contact into the direction of the buffer/absorber junction.

2. The minimisation of junction recombination, both at the point of equal capture rates of holes and electrons in the space charge region (recombination path B in Figure 15) as well as at the absorber/buffer interface (path C), requires an increased band gap towards the absorber surface. If one had the choice, one would clearly favour a decrease of the valence-band energy, as shown in Figure 15, over an increase of the conduction band energy. This favours a grading with the help of S/Se alloying, as at least a part of the increasing band-gap energy is accommodated by a decrease of the energy of the valence-band edge that should minimise interface recombination.

FIGURE 15 Band diagram of a ZnO/CdS/Cu(In,Ga)(Se,S)₂ heterojunction with a graded-gap absorber. The minimum band gap energy is in the quasi neutral part of the absorber. An increasing Ga/In ratio towards the back surface and an increasing Ga/In or S/Se-ratio towards the front minimise recombination in critical regions at the back contact (recombination path A’), in the space charge region (path B), and at the heterointerface (path C). The dotted lines correspond to the conduction and valence band edge energies of a non-graded device.

6 Conclusions

The objective of this chapter is the description of the recent achievements on Cu(In,Ga)Se₂-based solar cells as well as an account of our present understanding of the physical properties of the materials involved and the electronic behaviour of the devices. Cu(In,Ga)Se₂ is in a leading position among polycrystalline thin-film solar cell materials because of the benign, forgiving nature of the bulk materials and interfaces. Nevertheless, we want to guide the attention of the reader also in the direction of the work that has still to be done.

Three of the four cornerstones 1–4 for the recent achievements mentioned in Section 3.2 concern the growth of the films: the optimised deposition conditions and the incorporation of Na and Ga. However, no detailed model is available to describe the growth of Cu(In,Ga)Se₂, and especially the impact of Na, which, in our opinion, is the most important of the different ingredients available to tune the electronic properties of the absorber. A clearer understanding of Cu(In,Ga)Se₂ growth would allow us to find optimised conditions in the wide parameter space available, and thus to reduce the number of recombination centres and compensating donors and optimise the number of shallow acceptors. The deposition of the buffer layer, or more generally speaking, the formation of the heterojunction, is another critical issue. The surface chemistry taking place during heterojunction formation, and also during post-deposition treatments, is decisive for the final device performance. Both processes greatly affect not only the surface defects (i.e., recombination and charge), and therefore the charge distribution in the device, but also the defects in the bulk of the absorber. Concentrated effort and major progress in these tasks would not only allow us to push the best efficiencies further towards 20%, but would also provide a sound knowledge base for the various attempts at the commercialisation of Cu(In,Ga)Se₂ solar cells.

Acknowledgements

The authors thank all our colleagues at the IPE for various discussions and fruitful collaboration. We are especially grateful to J. H. Werner for his continuous support of the Cu(In,Ga)Se₂ research at the ipe. We are grateful to M. Turcu for a critical reading of the manuscript. Support of our work during many years by the German Ministry of Research, Technology, and Education (BMBF), by the Ministry of Economics, and by the European Commission is especially acknowledged.

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Table of Contents for Chapter IC-3. Cu(In,Ga)Se2 Thin-Film Solar Cells

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Chapter IC-3

Cu(In,Ga)Se2 Thin-Film Solar Cells

Chapter Outline

1 Introduction