CIGS introduction

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Cu(InGa)Se2 Solar Cells
William N. Shafarman1 and Lars Stolt2
 University of Delaware, Newark, DE, USA, 2 Uppsala University,
Uppsala, Sweden

       Cu(InGa)Se2 -based solar cells have often been touted as being among the most promising
       of solar cell technologies for cost-effective power generation. This is partly due to the
       advantages of thin films for low-cost, high-rate semiconductor deposition over large areas
       using layers only a few microns thick and for fabrication of monolithically interconnected
       modules. Perhaps more importantly, very high efficiencies have been demonstrated with
       Cu(InGa)Se2 at both the cell and the module levels. Currently, the highest solar cell effi-
       ciency is 18.8% with 0.5 cm2 total area fabricated by the National Renewable Energy
       Laboratory (NREL) [1]. Furthermore, several companies have demonstrated large area
       modules with efficiencies >12% including a confirmed 13.4% efficiency on a 3459 cm2
       module by Showa Shell [2]. Finally, Cu(InGa)Se2 solar cells and modules have shown
       excellent long-term stability [3] in outdoor testing. In addition to its potential advantages
       for large-area terrestrial applications, Cu(InGa)Se2 solar cells have shown high radia-
       tion resistance, compared to crystalline silicon solar cells [4, 5] and can be made very
       lightweight with flexible substrates, so they are also promising for space applications.
              The history of CuInSe2 solar cells starts with the work done at Bell Laboratories
       in the early 1970s, even though its synthesis and characterization were first reported by
       Hahn in 1953 [6] and, along with other ternary chalcopyrite materials, it had been char-
       acterized by several groups [7]. The Bell Labs group grew crystals of a wide selection of
       these materials and characterized their structural, electronic, and optical properties [7–9].
       The first CuInSe2 solar cells were made by evaporating n-type CdS onto p-type single
       crystals of CuInSe2 [10]. These devices were initially recognized for their potential as
       near-infrared photodetectors since their spectral response was broader and more uniform
       than Si photodetectors. Optimization for solar cells increased the efficiency to 12% as
       measured under outdoor illumination “on a clear day in New Jersey” [11].

Handbook of Photovoltaic Science and Engineering. Edited by A. Luque and S. Hegedus
 2003 John Wiley & Sons, Ltd ISBN: 0-471-49196-9
568       Cu(InGa)Se2 SOLAR CELLS

             There has been relatively little effort devoted to devices on single-crystal CuInSe2
      since this early work, in part because of the difficulty in growing high-quality crystals [12].
      Instead, nearly all the focus has gone to thin-film solar cells because of their inherent
      advantages. The first thin-film CuInSe2 /CdS devices were fabricated by Kazmerski et al.
      using films deposited by evaporation of CuInSe2 powder along with excess Se [13].
      However, thin-film CuInSe2 solar cells began to receive a lot of attention when the first
      high-efficiency, 9.4%, cells were demonstrated by Boeing [14]. At the same time, interest
      in Cu2 S/CdS thin-film solar cells waned owing to problems related to electrochemical
      instabilities and many of these researchers turned their focus to CuInSe2 .
              The Boeing devices were fabricated using CuInSe2 deposited by coevaporation,
      that is, evaporation from separate elemental sources [15], onto ceramic substrates coated
      with a Mo back electrode. Devices were completed with evaporated CdS or (CdZn)S
      deposited in two layers with undoped CdS followed by an In-doped CdS layer that served
      as the main current-carrying material [15]. Throughout the 1980s, Boeing and ARCO
      Solar began to address the difficult manufacturing issues related to scale-up, yield, and
      throughput leading to many advancements in CuInSe2 solar cell technology. The two
      groups pursued different approaches to CuInSe2 deposition, which today remain the most
      common deposition methods and produce the highest device and module efficiencies.
      Boeing focused on depositing Cu(InGa)Se2 by coevaporation, while ARCO Solar focused
      on a two-stage process of Cu and In deposition at a low temperature followed by a reactive
      anneal in H2 Se.
             The basic solar cell configuration implemented by Boeing provided the basis for a
      series of improvements that have lead to the high-efficiency device technology of today.
      The most important of these improvements to the technology include the following:

      • The absorber-layer band gap was increased from 1.02 eV for CuInSe2 to 1.1–1.2 eV by
        the partial substitution of In with Ga, leading to a substantial increase in efficiency [16].
      • The 1- to 2-µm-thick doped (CdZn)S layer was replaced with a thin, ≤50 nm, undoped
        CdS and a conductive ZnO current-carrying layer [17]. This increased the cell current
        by increasing the short wavelength (blue) response.
      • Soda lime glass replaced ceramic or borosilicate glass substrates. Initially, this change
        was made for the lower costs of the soda lime glass and its good thermal expansion
        match to CuInSe2 . However, it soon became clear that an increase in device performance
        and processing tolerance resulted primarily from the beneficial indiffusion of sodium
        from the glass [18].
      • Advanced absorber fabrication processes were developed that incorporate band gap
        gradients that improve the operating voltage and current collection [19, 20].

             From its earliest development, CuInSe2 was considered promising for solar cells
      because of its favorable electronic and optical properties including its direct band gap with
      high absorption coefficient and inherent p-type conductivity. As science and technology
      developed, it also became apparent that it is a very forgiving material since (1) high-
      efficiency devices can be made with a wide tolerance to variations in Cu(InGa)Se2
      composition [21, 22], (2) grain boundaries are inherently passive so even films with grain
      sizes less than 1 µm can be used, and (3) device behavior is insensitive to defects at
      the junction caused by a lattice mismatch or impurities between the Cu(InGa)Se2 and
                                                                     INTRODUCTION     569

CdS. The latter enables high-efficiency devices to be processed despite exposure of the
Cu(InGa)Se2 to air prior to junction formation.
       High-efficiency CuInSe2 -based solar cells have been fabricated by at least
10 groups around the world. While these groups employ a variety of processing
technologies, all the solar cells have the same basic cell structure built around
a Cu(InGa)Se2/CdS junction in a substrate configuration with a Mo back contact.
Figure 13.1 shows a cross-sectional schematic of a standard device. This structure utilizes
a soda lime glass substrate, coated with a sputtered Mo layer as a back contact. After
Cu(InGa)Se2 deposition, the junction is formed by chemical bath–deposited CdS with
thickness ≤50 nm. Then a high-resistance (HR) ZnO layer and a doped high-conductivity
ZnO layer are deposited, usually by sputtering or chemical vapor deposition. Either
a current-collecting grid or monolithic series interconnection completes the device or
module, respectively. A TEM micrograph of the same structure, shown in Figure 13.2,
clearly demonstrates the polycrystalline nature of these materials and the conformal
coverage of the CdS layer.

                                             collection grid

                                HR-Zno/n -ZnO (0.5 µm)

                                    CdS (0.05 µm)

                                  Cu(InGa)Se2 (2 µm)

                                     Mo (0.5 µm)

                                  Soda lime glass

           Figure 13.1 Schematic cross section of a typical Cu(InGa)Se2 solar cell





                 Figure 13.2 TEM cross section of a Cu(InGa)Se2 solar cell
570       Cu(InGa)Se2 SOLAR CELLS

             Several companies worldwide are pursuing the commercial development of
      Cu(InGa)Se2 -based modules. The most advanced, having demonstrated excellent
      reproducibility in its module manufacturing using the two-stage selenization process for
      Cu(InGa)(SeS)2 deposition [3], is Shell Solar Industries (SSI) in California, which was
      formerly ARCO Solar and then Siemens Solar. They are now in production with 5-, 10-,
      20-, and 40-W modules that are commercially available. In Germany, W¨ rth Solar is
      in pilot production using an in-line coevaporation process for Cu(InGa)Se2 deposition
      and has also reported large area modules with >12% efficiency. In the USA, several
      companies are in preproduction or pilot production: Energy Photovoltaics, Inc. (EPV) is
      using its own in-line evaporation process, International Solar Electric Technology (ISET)
      is developing a particle-based precursor for selenization, and Global Solar Energy (GSE)
      is pursuing a process for roll-to-roll coevaporation onto a flexible substrate. In Japan,
      Showa Shell, using a two-stage selenization process, and Matsushita, using coevaporation
      for Cu(InGa)Se2 deposition, are also in production development stages.
              Despite the level of effort on developing manufacturing processes, there remains
      a large discrepancy in efficiency between the laboratory-scale solar cells and minimod-
      ules, and the best full-scale modules. In part, this is due to the necessity for developing
      completely new processes and equipment for the large-area, high-throughput deposition
      needed for manufacturing thin-film photovoltaics. This is compounded by the lack of a
      comprehensive scientific base for Cu(InGa)Se2 materials and devices, due partly to the
      fact that it has not attracted a broader interest for other applications. This lack of a sci-
      ence base has been perhaps the biggest hindrance to the maturation of Cu(InGa)Se2 solar
      cell technology as most of the progress has been empirical. Still, in many areas a deeper
      understanding has emerged in the recent years.
             In this chapter we will review the current status and the understanding of thin-film
      Cu(InGa)Se2 solar cells from a technology perspective. For deeper scientific discussion of
      some aspects, we refer to suitable references. In order of presentation, this review covers
      (Section 13.2) structural, optical, and electrical properties of Cu(InGa)Se2 including a
      discussion of the influence of Na and O impurities; (Section 13.3) methods used to deposit
      Cu(InGa)Se2 thin films, the most common of which can be divided into two general types,
      multisource coevaporation and two-stage processes of precursor deposition followed by
      Se annealing; (Section 13.4) junction and device formation, which typically is done with
      chemical bath CdS deposition and a ZnO conduction layer; (Section 13.5) device operation
      with emphasis on the optical, current-collection, and recombination-loss mechanisms;
      (Section 13.6) module-manufacturing issues, including process and performance issues
      and a discussion of environmental concerns; and finally, (Section 13.7) a discussion of
      the outlook for CuInSe2 -based solar cells and critical issues for the future. In places where
      aspects of Cu(InGa)Se2 solar cells cannot be covered in full, reference will be made to
      other review works that serve to complement this chapter.

      The understanding of Cu(InGa)Se2 thin films, as used in photovoltaic (PV) devices, is
      primarily based on studies of its base material, pure CuInSe2 . Thorough reviews on
      CuInSe2 can be found in References [23–25]. However, the material used for making solar
      cells is Cu(InGa)Se2 containing significant amounts (of the order of 0.1%) of Na [26].
                                                                 MATERIAL PROPERTIES         571

     Even though the behavior of CuInSe2 provides a good basis for the understanding of
     device-quality material, there are pronounced differences when Ga and Na are present
     in the films. More recently, Cu(InGa)Se2 has been reviewed in the context of solar cells
     with an emphasis on electronic properties [27].
            In this section the structural, optical, and electrical properties of CuInSe2 are
     reviewed along with information about the surface and grain boundaries and the effect of
     the substrate. In each case, as appropriate, the effect of the alloying with CuGaSe2 to form
     Cu(InGa)Se2 and the impact of Na and O on the material properties will be discussed.

13.2.1 Structure and Composition
     CuInSe2 and CuGaSe2 have the chalcopyrite lattice structure. This is a diamond-like
     structure similar to the sphalerite structure but with an ordered substitution of the group I
     (Cu) and group III (In or Ga) elements on the group II (Zn) sites of sphalerite. This
     gives a tetragonal unit cell depicted in Figure 13.3 with a ratio of the tetragonal lattice
     parameters c/a close to 2 (see Table 13.1). The deviation from c/a = 2 is called the
     tetragonal distortion and stems from different strengths of the Cu–Se and the In–Se or
     Ga–Se bonds.
             The possible phases in the Cu–In–Se system are indicated in the ternary phase
     diagram in Figure 13.4. Thin films of Cu–In–Se prepared under an excess supply of Se,
     that is, normal conditions for thin-film growth of Cu(InGa)Se2 , have compositions that fall
     on, or close to, the tie-line between Cu2 Se and In2 Se3 . Chalcopyrite CuInSe2 is located on
     this line as well as a number of phases called ordered defect compounds (ODC), because
     they have a lattice structure described by the chalcopyrite structure with an ordered inser-
     tion of intrinsic defects. A comprehensive study of the Cu–In–Se phase diagram has been
     completed by G¨ decke et al. [32]. A detail of the Cu2 Se–In2 Se3 tie-line near CuInSe2 is
     described by the pseudobinary phase diagram reproduced in Figure 13.5 [32] Here α is
     the chalcopyrite CuInSe2 , δ is a high-temperature (HT) phase with the sphalerite structure,
     and β is an ODC phase. It is interesting to note that the single phase field for CuInSe2 at
     low temperatures is relatively narrow as compared to earlier beliefs, and does not contain




                     Figure 13.3   The unit cell of the chalcopyrite lattice structure
572       Cu(InGa)Se2 SOLAR CELLS

                                Table 13.1 Selected properties of CuInSe2
                         Property                                    Value                  Units   Reference
                                               a                       5.78                   A˚      [28]
      Lattice constant
                                               c                      11.62                   A˚
      Density                                                          5.75                 g/cm3     [28]
      Melting temperature                                              986                    C       [29]
      Thermal expansion                     (a axis)               8.32 × 10−6               1/K      [30]
        coefficients at 273 K                (c axis)               7.89 × 10−6               1/K
      Thermal conductivity                                            0.086                           [30]
        at 273 K
      Dielectric constant            Low frequency                 13.6 ± 2.4                         [31]
                                     High frequency                 8.1 ± 1.4
                                     Electrons                         0.09                           [30]
      Effective mass [me ]           Holes (heavy)                     0.71                           [30]
                                     Holes (light)                    0.092                           [30]
      Energy gap                                                      1.02                   eV       [30]
      Energy gap temperature                                       −2 × 10−4                eV/K      [30]


                                            CuSe2          ODC
                                                          phases      In2Se3
                                       CuSe                             InSe
                                    Cu2Se                                      In2Se

                               Cu                                                      In

      Figure 13.4 Ternary phase diagram of the Cu–In–Se system. Thin-film composition is usually
      near the pseudobinary Cu2 Se–In2 Se3 tie-line

      the composition 25% Cu. At higher temperatures, around 500◦ C, where thin films are
      grown, the phase field widens toward the In-rich side. Typical average compositions of
      device-quality films have 22 to 24 at.% Cu, which fall within the single-phase region at
      growth temperature.
              CuInSe2 can be alloyed in any proportion with CuGaSe2 , thus forming
      Cu(InGa)Se2 . Similarly, the binary phase In2 Se3 at the end point of the pseudobinary
      tie-line can be alloyed to form (InGa)2 Se3 , although it undergoes a structural change at
      Ga/(In + Ga) = 0.6 [33]. In high-performance devices, Ga/(In + Ga) ratios are typically
      0.2 to 0.3.
             One of the central characteristics of Cu(InGa)Se2 is its ability to accommodate large
      variations in composition without appreciable differences in optoelectronic properties.
                                                                           MATERIAL PROPERTIES   573

                                      800                                        785
                                      700                                  a+
                                                                           Cu2Se (HT)
                                      600             a +d       a


                                      500     b


                                                      a +b                 a+
                                                                           Cu2Se (RT)
                                         15           20              25                30

Figure 13.5 Pseudobinary In2 Se3 –Cu2 Se equilibrium phase diagram for compositions around the
CuInSe2 chalcopyrite phase, denoted α. The δ phase is the high-temperature sphalerite phase,
and the β phase is an ordered defect phase (ODC). Cu2 Se exists as a room-temperature (RT) or
high-temperature (HT) phase. (After G¨ decke T, Haalboom T, Ernst F, Z. Metallkd. 91, 622–634
(2000) [32])

This tolerance is one of the cornerstones of the potential of Cu(InGa)Se2 as a material
for efficient low-cost PV modules. Solar cells with high performance can be fabricated
with Cu/(In + Ga) ratios from 0.7 to nearly 1.0. This property can be understood from
theoretical calculations that show that the defect complex 2VCu + InCu , that is, two Cu
vacancies with an In on Cu antisite defect, has very low formation energy, and also that it
is expected to be electrically inactive [34]. Thus, the creation of such defect complexes can
compensate for Cu-poor/In-rich compositions of CuInSe2 without adverse effects on the
photovoltaic performance. Furthermore, crystallographic ordering of this defect complex is
predicted [34], which explains the observed ODC phases Cu2 In4 Se7 , CuIn3 Se5 , CuIn5 Se8 ,
and so on.
        The chalcopyrite phase field is increased by the addition of Ga or Na [35]. This can
be explained by a reduced tendency to form the ordered defect compounds owing to higher
formation energy for GaCu (in CuGaSe2 ) than for InCu (in CuInSe2 ). This leads to desta-
bilization of the 2VCu + InCu defect cluster related to the ODC phases [36, 37]. The effect
of Na in the CuInSe2 structure has been calculated by Wei et al. [38], with the result that
Na replaces InCu antisite defects, reducing the density of compensating donors. This the-
oretical result is supported by measurements of epitaxial Cu(InGa)Se2 films in which Na
is found to strongly reduce the concentration of compensating donors [37]. Together with
a tendency for Na to occupy Cu vacancies, the reduced tendency to form antisite defects
also suppresses the formation of the ordered defect compounds. The calculated effect of
Na is therefore consistent with the experimental observations of increased compositional
range in which single-phase chalcopyrite exists and increased conductivity [38, 39].
 574      Cu(InGa)Se2 SOLAR CELLS

13.2.2 Optical Properties
       The absorption coefficient α for CuInSe2 is very high, larger than 105 /cm for 1.4 eV and
       higher photon energies [40]. In many studies it was found that the fundamental absorption
       edge is well described by [30]

                                          α = A(E − Eg )2 /E                                (13.1)

       as for a typical direct band gap semiconductor. The proportionality constant A depends
       on the density of states associated with the photon absorption. From this relation, a band
       gap value of Eg = 1.02 ± 0.02 eV is obtained. The temperature dependence follows

                                    Eg (T ) = Eg (0) − a T 2 /(b + T )                      (13.2)

       where a and b are constants that vary between different measurements. In general, dEg /dT
       is about −2 × 10−4 eV/K [41].
              A rather complete picture of the optical properties of CuInSe2 and other Cu-ternary
       chalcopyrites is given in Reference [42]. Ellipsometric measurements of carefully pre-
       pared single-crystal samples were carried out and the dielectric functions were obtained
       together with the complex refractive index for different polarizations. From these mea-
       surements a band gap value for CuInSe2 of 1.04 eV was determined.
               A similar study was also made on bulk polycrystalline ingots of Cu(InGa)Se2
       having different compositions from x ≡ Ga/(Ga + In) = 0 to 1 [43]. Curves describing
       the complex refractive index, n + ik, for samples with x = 0 and 0.2 are reproduced in
       Figure 13.6. The complex refractive index can be used to calculate other optical param-
       eters like the absorption coefficient

                                               α = 4πk/λ                                    (13.3)

       In the same work the fundamental transitions for the different compositions were fit to
       an equation describing the band gap for CuIn1−x Gax Se2 as

                                Eg = 1.010 + 0.626x − 0.167x(1 − x)                         (13.4)

       In this equation the so-called bowing coefficient is 0.167. A value of 0.21 was obtained
       by theoretical calculations as compared to values in the range of 0.11 to 0.26 determined
       in various experiments [44].

13.2.3 Electrical Properties
       CuInSe2 with an excess of Cu is always p-type but In-rich films can be made p-type or
       n-type [45]. By annealing in a selenium overpressure, n-type material can be converted to
       p-type, and conversely, by annealing in a low selenium pressure, p-type material becomes
       n-type [46]. It is believed that this affects the concentration of Se vacancies, VSe , which
       act as compensating donors in p-type films. Device-quality Cu(InGa)Se2 films, grown
       with the excess Se available, are p-type with a carrier concentration of about 1016 /cm3 .
                                                              MATERIAL PROPERTIES         575


                                                           x = 0.2








                                        x = 0.2

                                 1      2            3         4       5

Figure 13.6 Complex refractive index for CuInSe2 and CuIn1−x Gax Se2 with x = 0.2 (After
Alonso M et al., Appl. Phys. A 74, 659–664 (2002) [43])

There is a large spread in mobility values reported for CuInSe2 . The highest values of hole
mobilities have been obtained for epitaxial films, where 200 cm2 /Vs has been measured for
Cu(InGa)Se2 with about 1017 /cm3 in hole concentration [37]. Single crystals have yielded
values in the range of 15 to 150 cm2 /Vs. Electron mobilities determined from single
crystals range from 90 to 900 cm2 /Vs [46]. Conductivity and Hall effect measurements
of thin-film samples are made cross-grain, but for device operation through-the-grain
values are more relevant, since individual grains may extend from the back contact to
the interface of the junction. The sheet conductivities of polycrystalline p-type films
correspond to mobility values of 5 to 50 cm2 /Vs, but it is quite possible that they are
limited by transport across grain boundaries.
        A large number of intrinsic defects are possible in the chalcopyrite structure.
Accordingly, a number of electronic transitions have been observed by methods such
as photoluminescence, photoconductivity, photovoltage, optical absorption, and electrical
measurements (see, for example, Reference [31]). However, it is difficult to assign transi-
tions to specific defects on an experimental basis. Instead, theoretical calculations of the
transition energies and formation energies provide a basis for identification of the different
intrinsic defects that are active in Cu(InGa)Se2 . Calculations of intrinsic defects in CuInSe2
and comparison with experimental data can be found in the comprehensive paper by Zhang
et al. [47]. A summary of their results is schematically shown in Figure 13.7. The defects
that are considered most important in device-quality material are presented in Table 13.2.
576       Cu(InGa)Se2 SOLAR CELLS


                                                 Cui(0/+)               D3
                                           0.8   InCu(0/+)
                                                 VIn(3−/2−)             D4
                                 Energy    0.6   CuIn(2−/−)
                                           0.4   VIn(2−/−)              A5

                                                 CuIn(−/0)              A4
                                                 VIn(−/0)               A3
                                                 VCu(−/0)               A1
                                          VBM       Theory       Experiment

      Figure 13.7 Electronic levels of intrinsic defects in CuInSe2 . On the left side the theoretical values
      are presented and on the right side experimentally reported values are presented. The height of the
      histogram columns on the right side represents the spread in experimental data. (From Zhang S,
      Wei S, Zunger A, Katayama-Yoshida H, Phys. Rev. B 57, 9642–9656 (1998) [47])

                               Table 13.2 The most important intrinsic defects
                               for device-quality CuInSe2
                               Defect       Energy position          Type
                                VCu          EV + 0.03 eV       Shallow acceptor
                                InCu         EC − 0.25 eV     Compensating donor
                                 VSe                          Compensating donor
                                CuIn         EV + 0.29 eV     Recombination center

              The effect of Ga on the electronic and defect properties is discussed in Refer-
      ence [36]. In those calculations, acceptor levels did not differ very much between CuInSe2
      and CuGaSe2 , but the donor levels are deeper in the Ga-containing compound. This is con-
      sistent with observations of increased p-type conductivity at high Ga-concentrations [48].
      At typical device compositions, Ga/(Ga + In) < 0.3, any effect of increased Ga content
      on conductivity has not been verified.

13.2.4 The Surface and Grain Boundaries
      The surface morphology and grain structure are most commonly characterized by scan-
      ning electron microscopy (SEM), but transmission electron microscopy (TEM) and atomic
      force microscopy have also proved valuable. A typical SEM image is shown in Figure 13.8
      and a TEM cross-sectional image in Figure 13.2. In general, the films used in devices
                                                         MATERIAL PROPERTIES            577


Figure 13.8 Scanning electron microscopy image of a typical Cu(InGa)Se2 film deposited on a
Mo-coated glass substrate by coevaporation

have grain diameters on the order of 1 µm but the grain size and morphology can vary
greatly depending on fabrication method and conditions. A variety of defects including
twins, dislocations, and stacking faults have been observed [49–51].
        It has been shown by X-ray photoelectron spectroscopy (XPS) that the free sur-
faces of CuInSe2 films with slightly Cu-poor composition have a composition close to
CuIn3 Se5 [52], corresponding to one of the ordered defect phases. Many attempts have
been made to identify such a layer on top of the films without success. It merely seems
as if the composition gradually changes from the bulk to the surface of the films. It was
proposed by Herberholz et al. [35] that band bending induced by surface charges drives
electromigrating Cu into the bulk leaving the surface depleted of Cu. This depletion is
stopped when the composition is that of CuIn3 Se5 , since further depletion requires a struc-
tural change of the material. Electromigration of Cu in CuInSe2 has been demonstrated
and also correlated with type conversion of the chalcopyrite material [53].
       The band bending as well as the CuIn3 Se5 composition of CuInSe2 surfaces dis-
appears when the material is exposed to atmosphere for some time as oxides form on the
surface. The surface oxidation is enhanced by the presence of Na [39]. The surface com-
pounds after oxidation have been identified as In2 O3 , Ga2 O3 , SeOx , and Na2 CO3 [54]. A
review of the surface and the interface properties can be found in Reference [55].
       It has been common practice to posttreat Cu(InGa)Se2 devices in air at typically
200◦ C. When devices were fabricated using vacuum-evaporated CdS or (CdZn)S to form
the junction, these anneals were often done for several hours to optimize the device
performance [14, 56]. The main effect associated with oxygen is explained as passivation
of selenium “surface” vacancies on the grains [57]. The VSe at the grain boundaries
can act as a recombination center. The positive charge associated with these donor-type
defects reduces the effective hole concentration at the same time that the intergrain carrier
transport is impeded. When oxygen substitutes for the missing selenium, these negative
effects are canceled.
        The overall noted beneficial effect of the presence of Na on the PV performance
of Cu(InGa)Se2 thin films lacks a complete explanation. In Reference [58] it is proposed
that the catalytic effect of Na on oxidation, by enhanced dissociation of molecular oxygen
into atomic oxygen, makes the passivation of VSe on grain surfaces more effective. This
model is consistent with the observation that Na and O are predominantly found at the
grain boundaries rather than in the bulk of the grains in CuInSe2 thin films [59].
578      Cu(InGa)Se2 SOLAR CELLS

13.2.5 Substrate Effects
      The effects of the substrate on the properties of thin-film polycrystalline Cu(InGa)Se2
      can be classified into three categories: (1) thermal expansion, (2) chemical effects, and
      (3) surface influence on nucleation.
             It can be assumed that after growth, when the substrate and film are still at the
      growth temperature, the stress in the Cu(InGa)Se2 film is low. The cooling down from
      growth temperature imposes a temperature change of about 500◦ C, and if the thermal
      expansion of the substrate and Cu(InGa)Se2 is different stress will be built up in the film.
      The thermal expansion coefficient for CuInSe2 is around 9 × 10−6 /K in the temperature
      interval of interest, which is similar to that of soda lime glass. A CuInSe2 film deposited
      on a substrate with a lower thermal expansion coefficient, such as borosilicate glass, will
      be under increasing tensile stress during cooldown. Typically, such films exhibit voids
      and microcracks [50]. When the thermal expansion coefficient of the substrate is higher
      than that of the film material, like for polyimide, it will result in compressive stress in
      the thin-film material, which may lead to adhesion failures.
              The most important effect of the soda lime glass substrate on Cu(InGa)Se2 film
      growth is that it supplies sodium to the growing chalcopyrite material. It has been clearly
      shown that this effect is distinct from the thermal expansion match of soda lime glass [60].
      The sodium diffuses through the Mo back contact, which also means that it is important
      to control the properties of the Mo [61]. The resulting microstructure of Cu(InGa)Se2
      is clearly influenced by the presence of Na with larger grains and a higher degree of
      preferred orientation, with the (112) axis perpendicular to the substrate. An explanation
      for this effect when high concentrations of Na are present has been proposed by Wei
      et al. [38].
             There is a wide range of preferred orientation between different growth processes,
      in spite of similar device performance. One reason for this variation is most likely the
      different properties of the surfaces on which the chalcopyrite material nucleates. A com-
      parison between Cu(InGa)Se2 grown on normal Mo-coated substrates and directly on
      soda lime glass shows that a much more pronounced (112) orientation occurs on glass, in
      spite of no difference in the Na concentration, as measured in the films afterwards [18].
      Further, the preferred orientation of the Cu(InGa)Se2 film has been shown to be correlated
      to the orientation of the Mo film [62] or an (InGa)2 Se3 precursor layer [63].

      A wide variety of thin-film deposition methods has been used to deposit Cu(InGa)Se2
      thin films. To determine the most promising technique for the commercial manufac-
      ture of modules, the overriding criteria are that the deposition can be completed at low
      cost while maintaining high deposition or processing rate with high yield and repro-
      ducibility. Compositional uniformity over large areas is critical for high yield. Device
      considerations dictate that the Cu(InGa)Se2 layer should be at least 1 µm thick and
      that the relative compositions of the constituents are kept within the bounds determined
      by the phase diagram, as discussed in Section 13.2.1. For solar cell or module fabrication,
      the Cu(InGa)Se2 is most commonly deposited on a molybdenum-coated glass substrate,
                                                              DEPOSITION METHODS             579

     though other substrate materials including metal or plastic foils have also been used and
     may have processing advantages.
            The most promising deposition methods for the commercial manufacture of mod-
     ules can be divided into two general approaches that have both been used to demonstrate
     high device efficiencies and in pilot scale manufacturing. The first approach is vacuum
     coevaporation in which all the constituents, Cu, In, Ga, and Se, can be simultaneously
     delivered to a substrate heated to 400 to 600◦ C and the Cu(InGa)Se2 film is formed in
     a single growth process. This is usually achieved by thermal evaporation from elemental
     sources at temperatures greater than 1000◦ C for Cu, In, and Ga. The second approach
     is a two-step process that separates the delivery of the metals from the reaction to form
     device-quality films. Typically the Cu, Ga, and In are deposited using low-cost and low-
     temperature methods that facilitate uniform composition. Then the films are annealed in
     a Se atmosphere, also at 400 to 600◦ C. The reaction and anneal step often takes longer
     time than formation of films by coevaporation due to diffusion kinetics, but is amenable
     to batch processing. High process rate can be achieved by moving continuously through
     sequential process steps or with a batch process whereby longer deposition or reaction
     steps can be implemented by handling many substrates in parallel.

13.3.1 Substrates
     Soda lime glass, which is used in conventional windows, is the most common sub-
     strate material used for Cu(InGa)Se2 since it is available in large quantities at low
     cost and has been used to make the highest efficiency devices. Cu(InGa)Se2 depo-
     sition requires a substrate temperature (TSS ) of at least 350◦ C and the highest effi-
     ciency cells have been fabricated using films deposited at the maximum temperature,
     TSS ≈ 550◦ C, which the glass substrate can withstand without softening too much [64].
     The glass is electrically insulating and smooth, which enables monolithic integration
     into modules.
            The soda lime glass has a thermal expansion coefficient of 9 × 10−6 /K [64], which
     provides a good match to the Cu(InGa)Se2 films. The glass composition typically includes
     various oxides such as Na2 O, K2 O, and CaO. These provide sources of alkali impurities
     that diffuse into the Mo and Cu(InGa)Se2 films during processing [18], producing the
     beneficial effects discussed in Section 13.2. However, a process that provides a more
     controllable supply of Na than diffusion from the glass substrate is preferred. This can
     be achieved by blocking sodium from the substrate with a diffusion barrier such as SiOx
     or Al2 O3 . Then sodium can be directly provided to the Cu(InGa)Se2 growth process by
     depositing a sodium-containing precursor layer onto the Mo film [65, 66]. Commercially
     available soda lime glass may also contain significant structural defects that can adversely
     impact module production [67]. Borosilicate glass does not contain the alkali impurities
     and may have fewer structural imperfections but has a lower thermal expansion coefficient,
     4.6 × 106 /K [64], and is more expensive.
             Substrates such as metal or plastic foils have advantages over glass substrates owing
     to their light weight and flexibility, which will be discussed in Section 13.6. Cu(InGa)Se2
     devices have been demonstrated with different metal and high-temperature polyimide
     substrates [68, 69].
580      Cu(InGa)Se2 SOLAR CELLS

13.3.2 Back Contact
      The Mo back contact, used for all high-efficiency devices, is typically deposited by direct
      current (dc) sputtering. The thickness is determined by the resistance requirements that
      depend on the specific cell or module configuration. A film with thickness 1 µm will
      typically have a sheet resistance of 0.1 to 0.2 / , a factor of 2 to 4 higher resistivity
      than bulk Mo. Sputter deposition of the Mo layer requires careful control of the pressure
      to control stress in the film [70] and to prevent problems such as poor adhesion that it
      might cause. During Cu(InGa)Se2 deposition, a MoSe2 layer forms at the interface [71].
      Its properties are influenced by the Mo film with less MoSe2 forming on dense Mo,
      sputter-deposited under low pressures [51]. This interfacial layer does not necessarily
      degrade device performance. Metals other than Mo have been investigated with limited
      success [72].

13.3.3 Coevaporation of Cu(InGa)Se2
      The highest efficiency devices have been deposited by thermal coevaporation from ele-
      mental sources [73]. An illustration of a laboratory system for Cu(InGa)Se2 coevaporation
      is shown in Figure 13.9. The process uses line-of-sight delivery of the Cu, In, Ga, and
      Se from Knudsen-type effusion cells or open-boat sources to the heated substrate. While
      the evaporation temperatures for each metal will depend on the specific source design,
      typical ranges are 1300 to 1400◦ C for Cu, 1000 to 1100◦ C for In, 1150 to 1250◦ C for
      Ga, and 300 to 350◦ C for Se evaporation.
             The sticking coefficients of Cu, In, and Ga are very high, so the film composition
      and growth rate are determined simply by the flux distribution and effusion rate from
      each source. The composition of the final film tends to follow the pseudobinary tie-line
      between (InGa)2 Se3 and Cu2 Se (see Figure 13.4) according to the relative concentration
      of Cu compared to In and Ga. The relative concentrations of In and Ga determine the band
      gap of the film, according to equation (13.4), and the effusion rates can be varied over
      the course of a deposition to change the film composition through its thickness. Se has a

                                                             Heater and



                                        To pump

                   Figure 13.9 Configuration for multisource elemental coevaporation
                                                         DEPOSITION METHODS            581

much higher vapor pressure and lower sticking coefficient, so it is always evaporated in
excess of that needed in the final film. Insufficient Se can result in a loss of In and Ga in
the form of In2 Se or Ga2 Se [74].
       Different deposition variations, using elemental fluxes deliberately varied over time,
have been explored using coevaporation. Four different sequences that have been used to
fabricate devices with efficiencies greater than 16% are shown in Figure 13.10. In each
case, the targeted final composition is Cu-deficient with Cu/(In + Ga) = 0.8 − 0.9. The
total deposition time may vary from 10 to 90 min, depending on the effusion rates from
the sources. So, for a film thickness of 2 µm, typical deposition rates vary from 20 to
200 nm/min.
       The first process is the simplest stationary process in which all fluxes are constant
throughout the deposition process [75]. In most cases, however, the fluxes are varied using
what is referred to as the Boeing process in which the bulk of the film is grown with Cu-
rich overall composition so that it contains a Cux Se phase in addition to Cu(InGa)Se2 [15].
The fluxes are then adjusted to finish the deposition with In- and Ga-rich flux so that the
final film composition has the desired Cu-deficient composition. One modification of this
is the second process shown in Figure 13.10. This process was first implemented with
CuInSe2 films deposited on non-Na containing substrates at TSS = 450◦ C, producing films
with increased grain size and improved device performance. The effect of Cux Se as a flux
for enhanced grain growth at higher TSS was proposed by Klenk et al. [76]. However,
in devices containing Na and Ga and with TSS > 500◦ C, no difference was found in the
device performance using films with Cu-rich or uniform growth processes [75].
        The third process shown in Figure 13.10 is a sequential process in which the In and
Ga are deposited separately from the Cu. This was first proposed by Kessler et al. [77]
with the deposition of an (InGa)x Sey compound, followed by the deposition of Cu and
Se until the growing film reaches the desired composition. The layers interdiffuse to
form the Cu(InGa)Se2 film. A modification by Gabor et al. [78] allows the Cu delivery
to continue until the film has an overall Cu-rich composition. Then a third step is added
to the process in which In and Ga, again in the presence of excess Se, are evaporated to
bring the composition back to Cu-deficient. The metals interdiffuse, forming the ternary
chalcopyrite film. This process has been used to produce the highest efficiency devices [1].
The improved device performance has been attributed to a band gap gradient, which
results from the Ga concentration decreasing from the Mo back contact to the film’s free
surface [19], and to improved crystallinity of the films [79].
       The last process shown in Figure 13.10 is an in-line process in which the flux
distribution results from the substrate moving sequentially over the Cu, Ga, and In sources.
This was first simulated in a stationary evaporation system [80] and has subsequently been
implemented by several groups in pilot manufacturing systems (see Section 13.6).
        A reproducible coevaporation process requires good control of the elemental fluxes
from each evaporation source. While the evaporation rates from each source can be con-
trolled simply by the source temperature, this may not give good reproducibility, especially
for the Cu source that is at the highest temperature. Open-boat sources in particular will
not give reproducible evaporation rates. Consequently, direct in situ measurement of the
fluxes is often used to control the evaporation sources. Electron impact spectroscopy [15],
quadrupole mass spectroscopy [81], and atomic absorption spectroscopy [82] have all
582      Cu(InGa)Se2 SOLAR CELLS

                                                                    TSS   550

                           Relative flux

                                           0.2   In                       450

                                                 Ga                       350

                                                                    TSS   550
                                           0.3   Cu
                           Relative flux

                                           0.2                            450
                                                                    TSS   550
                           Relative flux


                                           0.2                            450   [C]

                                                 Ga                       350

                                                                    TSS   550
                           Relative flux


                                                       Cu                 450

      Figure 13.10 Relative metal fluxes and substrate temperature for different coevaporation pro-
      cesses. In all cases, a constant Se flux is also supplied

      been successfully implemented. Direct flux measurement may be critical in a manufactur-
      ing scale process, particularly if source depletion over long run times causes the relation
      between source temperature and effusion rate to vary over time. In addition, the process
      can be monitored by in situ film thickness measurement using a quartz crystal monitor, or
      optical spectroscopy or X-ray fluorescence of the growing film [83]. The latter has also
      been used to measure composition. When the process includes a transition from Cu-rich
      to Cu-poor composition near the end of the deposition, it can be monitored by a change
                                                               DEPOSITION METHODS             583

     in the temperature resulting from a change in the emissivity of the film [84] or by the
     infrared transmission [85].
            The primary advantage of elemental coevaporation for depositing Cu(InGa)Se2
     films is its considerable flexibility to choose the process specifics and to control film
     composition and band gap. As proof of this flexibility, high-efficiency devices have been
     demonstrated using many process variations. The primary disadvantage results from the
     difficulty in control, particularly of the Cu-evaporation source, and the resulting need for
     improved deposition, diagnostic, and control technology. A second disadvantage is the
     lack of commercially available equipment for large-area thermal evaporation.

13.3.4 Two-step Processes
     The second common approach to Cu(InGa)Se2 film formation, usually referred to as two-
     step processing or selenization, has many variations in both the precursor deposition and
     the Se reaction steps. This general approach was first demonstrated by Grindle et al. [86]
     who sputtered Cu/In layers and reacted them in hydrogen sulfide to form CuInS2 . This
     was first adapted to CuInSe2 by Chu et al. [87]. The highest-efficiency Cu(InGa)Se2 cell
     reported using the reaction in H2 Se is 16.2%, on the basis of the active area [88], but there
     has been less effort at optimizing laboratory-scale cell efficiencies than with coevaporated
     Cu(InGa)Se2 . Showa Shell and Shell Solar have successfully scaled up this process to
     pilot commercial production and have demonstrated large-area module efficiencies as high
     as 13.4% [2].
            The metal precursor is used to determine the final composition of the film and
     to ensure spatial uniformity. Sputtering is an attractive process because it is easily scal-
     able using commercially available deposition equipment and can provide good uniformity
     over large areas with high deposition rates. However, other processes may have lower
     cost. CuInSe2 has been formed using metal precursor layers deposited by electrodeposi-
     tion [89], thermal or electron beam evaporation [90], screen printing [91], and application
     of nanoparticles [92]. Precursors that include Se, such as stacked layers of Cu/In/Se [93]
     or binary selenides, have also been used as precursor materials in various sequences
     and combinations [94]. Electrodeposition [95, 96] of Cu, In, Ga, and Se is effectively
     just another option for precursor deposition since the films similarly require a selenium
     reaction step.
            The precursor films are typically reacted in either H2 Se or Se vapor at 400 to
     500◦ C for 30 to 60 min to form the best device quality material. Poor adhesion [89]
     and formation of a MoSe2 layer [97] at the Mo/CuInSe2 interface may limit the reaction
     time and temperature. Reaction in H2 Se has the advantage that it can be done at atmo-
     spheric pressure and can be precisely controlled, but the gas is highly toxic and requires
     special precautions for its use. The precursor films can also be reacted in a Se vapor,
     which might be obtained by thermal evaporation, to form the CuInSe2 film [98]. A third
     reaction approach is rapid thermal processing (RTP) of either elemental layers, including
     Se, [99, 100] or amorphous evaporated Cu–In–Se layers [101].
           The reaction chemistry and kinetics for the conversion of Cu–In precursors to
     CuInSe2 has been characterized by X-ray diffraction of time-progressive reactions [102]
     and by in situ differential scanning calorimetry [103]. The results of these experiments
584      Cu(InGa)Se2 SOLAR CELLS

      describe CuInSe2 formation as a sequence of reactions starting with the formation of
      Cu11 In9 and In liquid, which will contain a small concentration of dissolved Cu. These
      react with Se to form a series of binary compounds. The formation of CuInSe2 then
      follows from
                                2 InSe + Cu2 Se + Se → 2 CuInSe2

      with complete reaction in ∼15 min at 400◦ C. The reaction path was shown to be the
      same for the reaction of Cu/In layers in either H2 Se or elemental Se [104].
             The addition of Ga, regardless of the precursor deposition sequence, does not
      readily give a film with uniformly increased band gap. Instead, all Ga in the reacted
      film accumulates near the Mo forming a CuInSe2 /CuGaSe2 structure, so the resulting
      device behaves like CuInSe2 [105] and lacks the increased operating voltage and other
      benefits of a wider band gap discussed in Section 13.5.4. Nevertheless, Ga inclusion
      provides improved adhesion of the CuInSe2 film to the Mo back contact and greater
      device performance, possibly owing to an improved structure with fewer defects [105].
      The Ga and In can be effectively interdiffused, converting the films to uniform band
      gap, by annealing in an inert atmosphere for 1 h at 600◦ C [106]. This anneal, however,
      may be impractical for commercial processing, so films in the best devices have the
      band gap increased by the incorporation of S near the front surface, forming a graded
      Cu(InGa)(SeS)2 layer [20, 107] that can give enhanced operating voltage in devices.
              The primary advantages of two-step processes for Cu(InGa)Se2 deposition are the
      ability to utilize more standard and well-established techniques for the metal deposition
      and reaction and anneal steps and to compensate for long reaction times with a batch
      processing mode or RTP of Se-containing precursors. Composition and uniformity are
      controlled by the precursor deposition and can be measured between the two steps. The
      biggest drawback to these processes is the limited ability to control composition and
      increase band gap, which may limit device and module performance. Other difficulties
      that must be overcome include poor adhesion and the use of hydrogen selenide, which is
      hazardous and costly to handle.

13.3.5 Other Deposition Approaches
      CuInSe2 -based films have been deposited using a wide range of thin-film deposition
      methods, in addition to those discussed above, which have been proposed as potential
      low-cost alternatives for manufacturing. These include reactive sputtering [108], hybrid
      sputtering in which Cu, In, and Ga are sputtered while Se is evaporated [109], closed-
      space sublimation [110], chemical bath deposition (CBD) [111], laser evaporation [112],
      and spray pyrolosis [113]. Great effort was made to explore different thin-film deposition
      techniques before coevaporation and the two-step processes above became dominant.
      These methods are reviewed in Reference [25].

      The first experimental device that indicated the potential for CuInSe2 in high-performance
      solar cells was a heterojunction between a p-type single crystal of CuInSe2 and a thin
      film of n-type CdS [10, 11]. Consequently, in the early thin-film work the junction was
                                                JUNCTION AND DEVICE FORMATION                585

     formed by depositing CdS on the CuInSe2 films [114]. The device was further developed
     to contain an undoped layer of CdS, followed by CdS doped with In, both deposited
     by vacuum evaporation [14]. This defined the device structure (see Figure 13.1), which
     is basically the same as is commonly used today since the doped CdS is functionally a
     transparent conductor. A performance gain was achieved by alloying the CdS with ZnS
     to widen the band gap [15]. Further improvement of the performance was achieved when
     the doped CdS layer was replaced with doped ZnO [115, 116]. The undoped CdS layer
     adjacent to the Cu(InGa)Se2 film was reduced in thickness in order to maximize the
     optical transmission. Since ZnO has a wider band gap than CdS, more light is transmitted
     into the active part of the device, resulting in a current gain. A conformal and pinhole-free
     coating of this thin CdS layer is obtained by using chemical bath deposition to make the
     CdS buffer layer.

13.4.1 Chemical Bath Deposition
     Chemical bath deposition (CBD) of thin-film materials can be viewed as a chemical vapor
     deposition (CVD) in the liquid phase instead of the gas phase. It is also referred to as
     solution growth. The method has been used in particular for chalcogenide materials such
     as PbS [117], CdS [118], and CdSe [119]. A variety of precursor compounds or ions can
     be used to deposit a specific compound.
           Deposition of CdS buffer layers on Cu(InGa)Se2 is generally made in an alkaline
     aqueous solution (pH > 9) of the following three constituents:

     1. a cadmium salt; for example, CdSO4 , CdCl2 , CdI2 , Cd(CH3 COO)2
     2. a complexing agent; commonly NH3 (ammonia)
     3. a sulfur precursor; commonly SC(NH2 )2 (thiourea).

     The concentrations of the various components of the solution can be varied over a range
     and each laboratory tends to use its own specific recipe. One example of a recipe that is
     being used to fabricate state-of-the-art Cu(InGa)Se2 solar cells is

     1. 1.4 × 1/103 M CdI2 or CdSO4
     2. 1 M NH3
     3. 0.14 M SC(NH2 )2

     The Cu(InGa)Se2 film is immersed in a bath containing the solution and the deposition
     takes place in a few minutes at a temperature of 60 to 80◦ C. This can be done either
     by immersion in a room-temperature bath that subsequently is heated to the desired
     temperature or by preheating the solution. The reaction proceeds according to the formula

            Cd(NH3 )4 2+ + SC(NH2 )2 + 2 OH− → CdS + H2 NCN + 4 NH3 + 2 H2 O

     In practice, the chemical bath deposition is typically done in the laboratory with a very
     simple apparatus consisting of a hot plate with magnetic stirring, a beaker holding the
     solutions into which the substrate is immersed, and a thermocouple to measure bath
 586       Cu(InGa)Se2 SOLAR CELLS

                                             Sample holder

                                                 Plastic lid

                                                                Glass container




                                                                PTFE-coated magnet
                                                                Magnetic stirrer

               Figure 13.11 Typical laboratory apparatus for chemical bath deposition of CdS

       temperature. A typical arrangement, incorporating a water bath for more uniform temper-
       ature, is shown in Figure 13.11. Scale-up of the CBD process for manufacturing will be
       discussed in Section 13.6.1.
              The growth of CdS thin films by CBD occurs from ion by ion reaction or by
       clustering of colloidal particles. Depending on the bath condition, the resulting CdS lattice
       structure may be cubic, hexagonal, or a mixture [120]. Under typical conditions used for
       Cu(InGa)Se2 solar cells, the relatively thin CdS layers grow ion by ion, resulting in dense
       homogeneous films [121] with mixed cubic/hexagonal or predominantly hexagonal lattice
       structure [51, 122, 123]. The films consist of crystallites with a grain size of the order of
       tens of nanometers [122].
               Compositional deviation from stoichiometry is commonly observed. In particular,
       films tend to be sulfur-deficient and contain substantial amounts of oxygen [124, 125].
       In addition to oxygen, significant concentrations of hydrogen, carbon, and nitrogen have
       also been detected in device quality films [126]. The concentration of these impurities
       has been correlated to a reduction of the optical band gap and the amount of cubic CdS
       in relation to hexagonal CdS [127].

13.4.2 Interface Effects
       The interface between the Cu(InGa)Se2 and the CdS is characterized by pseudoepitaxial
       growth of the CdS and intermixing of the chemical species. Electronic band alignment
       will be discussed in Section 13.5.3. Transmission electron microscopy has shown that
       chemical bath–deposited CdS layers on Cu(InGa)Se2 films exhibit an epitaxial relation-
       ship at the interface with (112) chalcopyrite Cu(InGa)Se2 planes parallel to the (111)
       cubic or (002) hexagonal CdS planes [51, 123]. The lattice mismatch is very small for
       pure CuInSe2 with a (112) spacing of 0.334 nm as compared to a spacing of 0.336 nm
       for (111) cubic and (002) hexagonal CdS. In Cu(InGa)Se2 the lattice mismatch increases
       with the Ga content. CuIn0.7 Ga0.3 Se2 and CuIn0.5 Ga0.5 Se2 have (112) spacing of 0.331 nm
                                                             JUNCTION AND DEVICE FORMATION        587



                           Lattice constant

                                                       Cd1−x ZnxS

                                                 0.0   0.2     0.4         0.6   0.8   1.0

     Figure 13.12 The lattice spacing of the (112) planes of CuIn1−x Gax Se2 and the (111) cubic or the
     (002) hexagonal planes of Cd1−x Znx S. Empirical data from References [128] ((CdZn)S) and [129]
     (CuInSe2 , CuGaSe2 , and (Cu(InGa)Se2 ) are included

     and 0.328 nm, respectively. Figure 13.12 displays the (112) spacing for Cu(InGa)Se2 as a
     function of Ga/(In + Ga) ratio together with the (111)/(002) spacing of CdS–ZnS alloys.
            When Cu(InGa)Se2 films are immersed in the chemical bath for deposition of CdS,
     they are also subjected to chemical etching of the surface. In particular, native oxides are
     removed by the ammonia [130]. Thus, the CBD process cleans the Cu(InGa)Se2 surface
     and enables the epitaxial growth of the CdS buffer layer.
            In early single-crystal work, p –n homojunction diodes were fabricated by indif-
     fusion of Cd or Zn into p-type CuInSe2 [131, 132] at 200 to 450◦ C. Investigations of
     CuInSe2 /CdS interfaces did show interdiffusion of S and Se above 150◦ C and rapid Cd
     diffusion into CuInSe2 above 350◦ C [133]. More recently, intermixing of the constituents
     of the Cu(InGa)Se2 /CdS heterojunction has been observed even when the relatively low-
     temperature CBD process is used for growth of the CdS layer [134]. Investigations of
     the effect of a chemical bath without the thiourea showed an accumulation of Cd on the
     Cu(InGa)Se2 surface, possibly as CdSe [130]. Accumulation of Cd on the Cu(InGa)Se2
     surface was also observed in the initial stage of CdS growth in the complete chemi-
     cal bath [135]. The results were not conclusive on whether any interfacial compound is
     formed, but TEM investigations showed the presence of Cd up to 10 nm into the Cu-
     deficient surface region of the Cu(InGa)Se2 layer [123]. At the same time, a reduction
     of the Cu concentration was noted. An interpretation in which Cu+ is replaced with
     Cd2+ is proposed, on the basis of the very close ion radii of these ions, 0.96 and 0.97,
     respectively. XPS and secondary ion mass spectrometry (SIMS) profiles of Cu(InGa)Se2
     films and CuInSe2 single crystals exposed to chemical baths without thiourea also show
     evidence of indiffusion or electromigration of Cd [136].

13.4.3 Other Deposition Methods
     In the early days of Cu(InGa)Se2 research, vacuum evaporation of 2 to 3-µm-thick CdS
     was the standard method to fabricate the junction and 9.4% efficiency was obtained with
588      Cu(InGa)Se2 SOLAR CELLS

      pure CuInSe2 absorbers [14]. With evaporation it is difficult to nucleate and grow very thin
      continuous CdS layers such as those normally used in current state-of-the-art Cu(InGa)Se2
      devices, and the optical transmission of the window will be limited to energies less than
      the CdS band gap, 2.4 eV. Substrate temperatures of 150 to 200◦ C are used to obtain
      good optical and electrical properties of the evaporated CdS films. This is substantially
      higher than the substrate temperature used for chemical bath deposition. Improved device
      performance was achieved by alloying the evaporated CdS with ZnS [15]. Mixed (CdZn)S
      has a wider band gap, allowing increased optical transmission, and better lattice match to
      Cu(InGa)Se2 than CdS.
             The main drawback with vacuum evaporation is poor conformal coating resulting
      in nonuniform and incomplete coverage of the sometimes relatively rough Cu(InGa)Se2
      films. Sputter deposition leads to more conformal coverage. The general success of sput-
      tering for industrial large-area deposition motivated the exploration of sputter-deposited
      CdS buffer layers. Using optical emission spectroscopy to control the sputtering pro-
      cess, Cu(InGa)Se2 devices with efficiencies up to 12.1% were fabricated, as compared
      to 12.9% for reference cells with chemical bath–deposited CdS [137]. Both evapora-
      tion and sputtering are vacuum processes, which can be incorporated in-line with other
      vacuum processing steps and do not create any liquid wastes. Still, CBD remains the
      preferred process for the CdS layer owing to its advantages in forming thin conformal
              Atomic layer chemical vapor deposition (ALCVD) is a method that also allows
      accurate control of the growth of thin conformal layers [138]. The method is being indus-
      trially used for deposition of another II-VI compound, ZnS. Inorganic precursors for
      deposition of CdS require the substrate temperature to be excessively high (>300◦ C) and
      work with organic precursors has been limited. The strong driving force for replacement
      of the environmentally nondesirable cadmium has focused the development of ALCVD on
      materials other than CdS. This is also valid for regular CVD, although some metal organic
      CVD (MOCVD) work has been reported. The full potential for chemical vapor–deposited
      CdS has therefore not been explored.
            Electrodeposition can be used to deposit CdS films but its use has not been reported
      in Cu(InGa)Se2 devices.

13.4.4 Alternative Buffer Layers
      The cadmium content in Cu(InGa)Se2 PV modules with CBD CdS buffer layers is low.
      Investigations show that the cadmium in Cu(InGa)Se2 modules can be handled safely,
      both with respect to environmental concerns and hazards during manufacturing (see
      Section 13.6.5). In spite of this, it would be preferable to eliminate cadmium in new
      products. There are in principle two approaches to Cd-free devices: (1) finding a buffer
      material that replaces CdS and (2) omitting the CdS layer and depositing ZnO directly
      onto the Cu(InGa)Se2 film. In practice, the two approaches tend to merge when the chem-
      ical bath deposition of CdS is replaced with a surface treatment of the Cu(InGa)Se2 with
      no or negligible film deposition before the subsequent deposition of the ZnO.
             A number of approaches and materials have been tried. A selection of promising
      results are presented in Table 13.3.
                                                    JUNCTION AND DEVICE FORMATION              589

Table 13.3 Performance of Cu(InGa)Se2 thin-film solar cells with various buffer layers and junc-
tion-formation methods alternative to chemical bath deposition of CdS
Buffer                          Deposition method     Efficiency    VOC     JSC        FF   Reference
material                                                 [%]      [mV]   [mA/cm2 ]   [%]
None                                                    10.5a     398      39.0      68      [139]
None                                                    15.0b     604      36.2      69       [1]
ZnO                       MOCVD                         13.9a     581      34.5      69      [140]
ZnO                       ALCVD                         11.7      512      32.6      70      [141]
Zn treatment              ZnCl2 solution                14.2b     558      36.3      70      [142]
Zn(O,S,OH)x               Chemical bath                 14.2c,d   567e     36.6e     68      [143]
ZnS                       Chemical bath                 16.9a,b   647      35.2      74      [144]
Zn treatment + ZnS        Chemical bath + ILGARf        14.2      559      35.9      71      [145]
Zn(Se,OH)                 Chemical bath                 13.7b,d   535      36.1      71      [146]
ZnSe                      ALCVD                         11.6a     502      35.2      65      [147]
ZnSe                      MOCVD                         11.6      469      35.8      69      [148]
Inx Sey                   Coevaporation                 13.0a     595      30.4      72      [149]
ZnInx Sey                 Coevaporation                 15.1      652      30.4      76      [150]
Inx (OH,S)y               Chemical bath                 15.7a,b   594      35.5      75      [151]
In2 S3                    ALCVD                         13.5      604      30.6      73      [152]
a Active area
b With antireflection layer
c Minimodule
d Confirmed
e Recalculated to single-cell   values
f Ion   Layer Gas Reaction

        When the numbers in Table 13.3 are analyzed, one must keep in mind that the qual-
ity of the Cu(InGa)Se2 layer varies significantly between the experiments. For example, in
the early results with direct ZnO [139], the reference cells with chemical bath–deposited
CdS showed 12.4% efficiency, whereas the 15% efficiency results [1] are obtained from
Cu(InGa)Se2 , which at best yielded an efficiency of 18.8%. On the other hand, an inferior
junction-formation method may cause a larger degradation of cell efficiency at higher effi-
ciency levels, since its defects may be relatively more important to the cell performance.
In order to evaluate the various Cd-free junction-formation methods from that respect, the
efficiency from each experiment is displayed in Figure 13.13 together with its reference,
or estimation thereof. In most cases, the Cd-free device is comparable to the CBD–CdS
device within typical variations.
       Altogether, it appears as if there are several possibilities for obtaining high effi-
ciency without Cd. All the listed methods include one or more of the elements Zn, In, and
S. Zn is directly included in most of the buffer materials or indirectly as ZnO transparent
contact with Inx Sey , In(OH,S)x , and In2 Se3 . Indications that n-type doping with Zn occurs
similarly to that with Cd have been found by the treatment of Cu(InGa)Se2 in Cd and Zn
solutions [136], and are consistent with junction formation by solid-state diffusion into
single crystals [132].
       In Figure 13.13 a slight tendency can be noted toward larger difference between
Cd-free and CdS reference cells for the direct ZnO approaches. It appears as if a buffer
layer between the Cu(InGa)Se2 and the ZnO is beneficial. Such a layer could passivate the
590      Cu(InGa)Se2 SOLAR CELLS

                                                                                                                        Cd-free                                  CBD-CdS














      Figure 13.13 The efficiency of Cu(InGa)Se2 solar cells with a selection of Cd-free junction-
      formation methods together with corresponding values of Cu(InGa)Se2 cells with chemical
      bath–deposited CdS

      Cu(InGa)Se2 surface, which would reduce the recombination in a shallow n-type emitter,
      and possibly also serve to protect the junction and near-surface region during subsequent
      deposition of the transparent contact materials.

13.4.5 Transparent Contacts
      The early Cu(InGa)Se2 devices used CdS doped with In or Ga as front-contact layers in
      addition to the CdS buffer layer. Short wavelength light (<520 nm) was absorbed near
      the surface in the thick CdS layer and did not generate any photocurrent. When chemical
      bath deposition allowed CdS buffer layers to be thin enough such that it no longer limited
      the short wavelength collection in the Cu(InGa)Se2 , photocurrent could be gained by
      increasing the band gap of the contact layer. Since the contact layer must also have high
      conductivity for lateral current collection, the obvious choice is a transparent conducting
      oxide (TCO), a class of materials used in such devices as displays and low-emission
      coatings on window glass panes. There are three main materials in this class: SnO2 ,
      In2 O3 :Sn (ITO), and ZnO. SnO2 requires relatively high deposition temperatures that
      restrict the potential in Cu(InGa)Se2 devices that cannot withstand temperatures greater
      than 200 to 250◦ C after CdS is deposited. ITO and ZnO can both be used, but the most
      common material is ZnO, favored by potentially lower material costs. A good overview
      of TCO thin-film materials can be found in Reference [153].
             The most commonly used low-temperature deposition method for TCO films is
      sputtering. ITO layers are routinely fabricated on an industrial scale using dc sputtering.
      Industrial practice is to use ceramic ITO targets and to sputter in an Ar:O2 mixture.
      Typical sputter rates range between 0.1 to 10 nm/s, depending on the application [154].
               Sputtering of doped ZnO films is not as developed as is sputtering of ITO. Neverthe-
      less, it is the preferred method for depositing the transparent front contact on Cu(InGa)Se2
                                                JUNCTION AND DEVICE FORMATION                591

     devices, with and without CBD–CdS, in the majority of the R&D groups. Typically,
     ZnO:Al films are deposited by radio frequency (rf) magnetron sputtering from ceramic
     ZnO:Al2 O3 targets with 1 or 2 weight% Al2 O3 . In large-scale manufacturing, dc sputter-
     ing from ceramic targets is favored since it requires simpler equipment and offers higher
     deposition rates [155].
             Reactive dc sputtering from Al/Zn alloy targets has also been used in the fab-
     rication of Cu(InGa)Se2 /CdS devices with the same performance as with rf sputtered
     ZnO:Al [156]. The use of Zn/Al alloy targets allows lower costs than ceramic ZnO:Al2 O3
     targets, but reactive sputtering requires very precise process control owing to the so-called
     hysteresis effect [157] so that optimal optoelectronic properties are achieved only within
     a very narrow process window. Deposition rates in 4 to 5 nm/s range have been achieved.
            Chemical vapor deposition (CVD) provides another deposition option and is used
     by one commercial manufacturer of Cu(InGa)Se2 modules to deposit ZnO [158]. The
     reaction occurs at atmospheric pressure between water vapor and diethylzinc and the
     films are doped with fluorine or boron.
            ALCVD deposition of ZnO has also been tested [159]. The atomic layer by atomic
     layer growth gives very low deposition rates, but the surface-controlled growth process
     gives uniform layers within a wide process window concerning reactant flow. This allows
     large batches to be processed, resulting in a reasonable throughput in spite of the limited
     growth rate.
            As with the Mo back contact, the requirements for sheet resistance of the trans-
     parent contact layer will depend on the specific cell or module design. Typically, small
     area cells use layers with 20–30 / , while modules may require 5–10 / . In either
     case, the sheet resistance is usually controlled by the layer thickness.

13.4.6 Buffer Layers
     It is common practice to use a buffer layer of undoped high-resistivity (HR) ZnO before
     sputter deposition of the TCO layer. Depending on the deposition method and conditions,
     this layer may have a resistivity of 1–100 cm compared to the transparent contact with
     10−4 –10−3 cm. Typically, 50 nm of HR ZnO is deposited by rf magnetron sputtering
     from an oxide target.
            The gain in performance by using an HR ZnO buffer layer in ordinary devices
     with CBD–CdS is related to the CdS thickness [156, 160, 161]. One explanation of
     the role of a ZnO buffer layer is given by [160] as resulting from locally nonuniform
     electronic quality of the Cu(InGa)Se2 layer that can be modeled by a parallel diode with
     high recombination current. The influence of these regions on the overall performance is
     reduced by the series resistance of the HR ZnO layer. This series resistance has a negligible
     effect on the performance of the dominant parts of the device area. A related explanation
     would attribute the local areas with poor diode characteristics caused by pinholes in the
     CdS layer, which create parallel diodes with a Cu(InGa)Se2 /ZnO junction. In this case
     improved diode quality due to the ZnO buffer would improve overall performance. Either
     case is consistent with the observation that a beneficial effect from the ZnO buffer is not
     observed when the CBD–CdS layer is thick enough [161].
592       Cu(InGa)Se2 SOLAR CELLS

             Another potential reason for using an HR ZnO buffer layer is to add protection
      of the interface region from sputter damage induced during deposition of the TCO layer
      which typically requires more harsh conditions. This seems to be particularly important for
      some alternative Cd-free buffer layers or with dc magnetron–sputtered TCO layers [162].

13.4.7 Device Completion
      In order to contact laboratory test cells, a metal contact is deposited onto the TCO layer. It
      is shaped as a grid with minimum shadow area in order to allow as much light as possible
      into the device. Solar cell measurement standards recommend a minimum cell area of
      1 cm2 , but many labs routinely use cells in the order of 0.5 cm2 . The metal grid contact
      can be made by first depositing some tens of nanometers of Ni to prevent the formation
      of a high resistance oxide layer, and subsequently depositing a few micrometers of Al.
      Evaporation through an aperture mask is a suitable deposition method.
             After deposition of the metal grid, the total cell area is defined by removing the
      layers on top of the Mo outside the cell area by mechanical scribing or laser patterning.
      Alternatively, just the layers on top of the Cu(InGa)Se2 can be removed, by photolitho-
      graphy and etching, since the lateral resistance of the Cu(InGa)Se2 prevents collection
      outside the cell area.
             The only significant difference in the device layers between lab cells and modules
      is the thickness of the TCO. Modules normally do not have any grid that assists in
      current collection over the cell area, so a substantially thicker TCO layer, that is, higher
      sheet conductivity, is needed in order to keep resistive losses low. A TCO layer with
      higher sheet conductivity may also have lower optical transmission in the infrared due to
      increased free-carrier absorption resulting in a decreased photocurrent.

      Cu(InGa)Se2 solar cells have achieved efficiencies approaching 20%, the highest of any
      thin-film solar cells, largely by empirical processing improvements and in spite of rela-
      tively poor understanding of the underlying mechanisms and electronic defects that control
      the device behavior. However, a more complete picture of the device operation is emerg-
      ing to enable both a better understanding of the devices and identification of pathways to
      further improvements.
             The operation of Cu(InGa)Se2 /CdS solar cells is characterized by high quantum
      efficiency (QE ) and short-circuit current. The open-circuit voltage increases with the
      band gap of the absorber layer and is insensitive to grain boundaries and defects at
      the Cu(InGa)Se2 /CdS interface. A basic device model can be constructed in which the
      voltage is limited by recombination through bulk trap states in the space charge region of
      the Cu(InGa)Se2 absorber layer. Recombination at the Cu(InGa)Se2 /CdS interface is min-
      imized by proper doping and band alignment or surface treatment to create an effective
      n-type inversion layer in the near-junction region of the absorber layer.
             The device operation can be described by identifying loss mechanisms. These
      can be divided into three categories. The first are optical losses that limit generation of
      carriers and therefore the device current. The second are recombination losses that limit
                                                                                     DEVICE OPERATION   593

     the voltage. Finally, there are parasitic losses, such as series resistance, shunt conductance,
     and voltage-dependent current collection, which are most evident by their effect on the
     fill factor but can also reduce JSC and VOC .

13.5.1 Light-generated Current
     The highest efficiency Cu(InGa)Se2 device has JSC = 35.2 mA/cm2 [1] out of a possible
     42.8 mA/cm2 available for a band gap of 1.12 eV under AM1.5 global illumination. Quan-
     tum efficiency is a valuable tool to characterize the losses responsible for this difference in
     current. The light-generated current is the integral of the product of the external quantum
     efficiency (QE ext ) and the illumination spectrum. QE ext is controlled by the band gap of
     the Cu(InGa)Se2 absorber layer, the CdS and ZnO window layers, and a series of loss
     mechanisms. These losses are illustrated in Figure 13.14 where typical QE curves at two
     different voltage biases, 0 V and −1 V, are shown. The QE curve at −1 V is slightly
     higher at longer wavelengths. The current loss under 100 mW/cm2 illumination is listed
     in Table 13.4 for each of these mechanisms. Losses 1 to 5 are optical and 6 is electronic.
     In practice, the magnitude of each of these losses will depend on the details of the device
     design and optical properties of the specific layers. The losses include the following:

     1. Shading from a collection grid used for most devices. In an interconnected module this
        will be replaced by the area used for the interconnect, as discussed in Section 13.6.2.
     2. Front surface reflection. On the highest-efficiency devices this is minimized with an
        antireflection layer for which an evaporated MgF2 layer with thickness ∼100 nm is
        commonly used. However, this is not practical in a module in which a cover glass is
        typically required.

                                                    3.0    2.0        1.5                1.0

                                              0.8                                              (2)
                         Quantum efficiency

                                              0.6                     (5)



                                                    400   600       800       1000     1200

     Figure 13.14 Quantum efficiency (solid lines) at 0 V and −1 V and optical losses for a
     Cu(InGa)Se2 /CdS solar cell in which the Cu(InGa)Se2 has Eg = 1.12 eV
594      Cu(InGa)Se2 SOLAR CELLS

                 Table 13.4 Current loss, J , for E > 1.12 eV due to the optical and
                 collection losses illustrated in Figure 13.14 for a typical Cu(InGa)Se2 /CdS
                 solar cell
                 Region in                  Optical loss mechanism                    J
                 Figure 13.14                                                     [mA/cm2 ]
                      (1)          Shading from grid with 4% area coverage            1.7
                      (2)          Reflection from Cu(InGa)Se2 /CdS/ZnO                3.8
                      (3)          Absorption in ZnO                                  1.8
                      (4)          Absorption in CdS                                  0.8
                      (5)          Incomplete generation in Cu(InGa)Se2               1.9
                      (6)          Incomplete collection in Cu(InGa)Se2               0.4

      3. Absorption in the TCO layer. Typically, there is 1 to 3% absorption through the visible
         wavelengths, which increases in the near IR region, λ > 900 nm, where free-carrier
         absorption becomes significant, and for λ < 400 nm near the ZnO band gap.
      4. Absorption in the CdS layer. This becomes appreciable at wavelengths below ∼520 nm
         corresponding to the CdS band gap 2.42 eV. The loss in QE for λ < 500 nm is
         proportional to the CdS thickness since it is commonly assumed that electron–hole
         pairs generated in the CdS are not collected. Figure 13.14 shows a device with a
         ∼30 nm-thick CdS layer. In practice, the CdS layer is often thicker and the absorption
         loss greater.
      5. Incomplete absorption in the Cu(InGa)Se2 layer near the Cu(InGa)Se2 band gap. Band
         gap gradients, resulting from composition gradients in many Cu(InGa)Se2 films, also
         affect the steepness of the long-wavelength part of the QE curve. If the Cu(InGa)Se2 is
         made thinner than ∼1.0 µm, this loss becomes significant [163] because of insufficient
         absorption at long wavelengths.
      6. Incomplete collection of photogenerated carriers in the Cu(InGa)Se2, discussed below.

            QE ext is then given by

             QE ext (λ, V ) = [1 − R(λ)][1 − AZnO (λ)][1 − ACdS (λ)] QE int (λ, V )             (13.5)

      where R is the total reflection, including the grid shading, AZnO is the absorption in
      the ZnO layer and ACdS is the absorption in the CdS layer. QE int , the internal quantum
      efficiency, is the ratio of photogenerated carriers collected to the photon flux that arrives
      at the absorber layer and can be approximated by [164]

                                                      exp[−α(λ)W (V )]
                                 QE int (λ, V ) ∼ 1 −
                                                =                                               (13.6)
                                                          αL + 1

      where α is the Cu(InGa)Se2 absorption coefficient, W is the space charge width in the
      Cu(InGa)Se2 , and L is the minority carrier diffusion length. This approximation assumes
      that all carriers generated in the space charge region are collected without recombination
      loss. Since W is a function of the applied voltage bias, QE int and total light-generated
      current are, in general, voltage-dependent, so the latter can be written as JL (V ). Values
      of W in the range 0.1–0.5 µm have been reported for typical cells at 0 V.
                                                                                   DEVICE OPERATION    595

                                        0.8           800


                                                                             1100 nm


                                          0.0               0.5              1.0        1.5

       Figure 13.15 Absorption of light with different wavelengths in Cu(InGa)Se2 with x = 0.2

            The absorption of light with different wavelengths in Cu(InGa)Se2 with x = 0.2
     is shown in Figure 13.15. At thickness d, this is given by exp(−αd) with α calculated
     at each wavelength using equation (13.3) and the data in Figure 13.6. If the effective
     collection length L + W is smaller than 0.5 to 1 µm, a significant fraction of electrons
     are generated deeper into the Cu(InGa)Se2 layer, and their incomplete collection can be a
     significant loss mechanism for Cu(InGa)Se2 devices [116, 165]. The effect of JL (V ) on
     current–voltage behavior increases with forward voltage bias and therefore has its largest
     effect on the fill factor and VOC [166, 167]. The effect of a voltage-dependent collection
     on JSC is illustrated in Figure 13.14 by the increase in QE measured at −1 V applied
     voltage bias compared to that measured at 0 V.

13.5.2 Recombination
     The current–voltage (J –V ) behavior of Cu(InGa)Se2 /CdS devices can be described by a
     general diode equation:
                      J = JD − JL = JO exp                            (V − RS J ) + GV − JL           (13.7)
     with the diode current JO given by:

                                                JO = JOO exp −                                        (13.8)

     The ideality factor A, barrier height b , and prefactor JOO depend on the specific recombi-
     nation mechanism that dominates JO , while the series resistance RS and shunt conductance
     G are losses that occur in series or parallel with the primary diode. General expres-
     sions for A, b , and JOO in the cases of recombination through the interface, space
     charge region, or bulk of the absorber layer can be found in various textbooks (see, for
     example [168]).
596       Cu(InGa)Se2 SOLAR CELLS

                                                                        x      Eg [eV]
                                                                  a     0        1.02
                                                                  b     0.24     1.16
                                                                  c     0.61     1.40

                            Current                                                             a    b     c


                                                        −0.3             0.0             0.3        0.6        0.9

      Figure 13.16 Current–voltage curves for Cu(InGa)Se2 /CdS solar cells with different relative Ga
      content giving (a) Eg = 1.02, (b) 1.16, and (c) 1.4 eV

                                                            3.0        2.0        1.5                1.0

                             Quantum efficiency


                                                                                   c            b    a


                                                        400           600       800        1000     1200   1400

              Figure 13.17 Quantum efficiency curves for the devices shown in Figure 13.16

             To understand the specific diode behavior of Cu(InGa)Se2 /CdS solar cells, it is
      instructive to look at the effect of the Cu(InGa)Se2 band gap, varied by changing x ≡
      Ga/(In + Ga), and temperature. Figures 13.16 and 13.17 show J –V and QE curves for
      3 devices with x = 0, 0.24, and 0.61, corresponding to Eg = 1.02, 1.16, and 1.40 eV,
      respectively. VOC increases and the position of the long-wavelength QE edge shifts to
      greater energy as Eg increases. Figure 13.18 shows the temperature dependence of VOC for
      these devices. In each case, as T → 0, VOC → Eg /q. Thus, combining equations (13.7)
                                                                     DEVICE OPERATION    597




                             0.8               a



                                   0               100         200      300

     Figure 13.18 Temperature dependence of VOC for the devices shown in Figure 13.16

and (13.8) and assuming G               JL /VOC , the open-circuit voltage becomes

                                               Eg   AkT    JOO
                                       VOC =      −     ln                              (13.9)
                                               q     q      JL

with the barrier height      b   = Eg .
        The different recombination paths are effectively connected in parallel so that
VOC will be controlled by the single dominant mechanism with the highest current.
The values of b and A can be used to distinguish between recombination in the bulk
absorber, in the space charge region of the Cu(InGa)Se2 , or at the Cu(InGa)Se2 /CdS
interface [27, 169]. Each of the curves in Figure 13.16 can be fit to equation (13.7) with
A = 1.5 ± 0.3. For a wide range of thin-film solar cells, it has been demonstrated that
VOC (T → 0) = QE g and 1 < A < 2 similar to the data above. Specifically, this has been
shown for CuInSe2 [116, 170] and Cu(InGa)(SeS)2 [171] devices, independent of the
(CdZn)S buffer-layer band gap [170], and for a variety of different absorber-layer depo-
sition processes [172]. These results for b and A indicate that Cu(InGa)Se2 /CdS solar
cells operate with the diode current controlled by Shockley–Read–Hall type recombi-
nation in the Cu(InGa)Se2 layer [168]. This recombination is greatest through deep trap
states in the space charge region of the Cu(InGa)Se2 where there are comparable sup-
plies of electrons and holes available, that is, p ≈ n. The variation in A between 1 and
2 depends on the energies of the deep defects that act as dominant trap states [173]. As
these states move toward the band edges, A → 1 and the recombination becomes closer
to band-to-band bulk recombination.
      These observations exclude recombination in the neutral bulk region of the
absorber layer, which should give A = 1. Interface recombination would give b <
Eg [Cu(InGa)Se2 ] with a dependence on the (CdZn)S band gap, although A might vary
from 1 to >2 [174]. Back surface recombination at the Mo/Cu(InGa)Se2 interface will
598       Cu(InGa)Se2 SOLAR CELLS

      be negligible so long as the minority-carrier diffusion length is small compared to the
      total Cu(InGa)Se2 thickness. If L + W ≈ d, a back surface field may be implemented,
      for example, by increasing the Ga content near the Mo to give a band gap gradient.
             In real Cu(InGa)Se2 materials with imperfect structures, trap defects will not exist at
      discrete energies but form defect bands or tails at the valence and conduction bands. Then
      the total recombination current can be determined by integrating over the defect spec-
      trum. Recombination through an exponential bandtail was used to explain the temperature
      dependence in A observed in some devices [175]. Analysis of the temperature dependence
      of A was further explained by a tunneling enhancement of the recombination current,
      particularly at reduced temperatures [176]. The same defects in the Cu(InGa)Se2 space
      charge region that control recombination were also used to explain observed metastabilities
      including persistent photoconductivity and open-circuit voltage decay [177]. Admittance
      spectroscopy has proved to be a useful tool to characterize the distribution of electronic
      defects in Cu(InGa)Se2 /CdS solar cells [178] and the density of an acceptor state ∼0.3 eV
      from the valence band has been correlated to VOC [179]. The minority-carrier lifetime is
      another valuable parameter to characterize Cu(InGa)Se2 /CdS devices. Transient photocur-
      rent [180] and time-resolved photoluminescence [167] measurements each were used to
      calculate lifetimes in the range of 10 to 100 ns for high-efficiency devices. Still, a critical
      problem that remains is to identify which of the calculated or measured defects discussed
      in Section 13.2 provides for the recombination traps that limit voltage in the devices. A
      good review of the characterization of electronic defects and their effect on Cu(InGa)Se2
      devices is provided by Rau and Schock [27].
              In practice, analysis of J –V data is commonly used to determine the diode param-
      eters JO , A, and b . This requires that RS and G are negligible, or suitable corrections are
      made to the data, and that JL is independent of V . Failure to account for JL (V ) can lead to
      errors in analysis of current–voltage data [166] and in many cases the fundamental diode
      parameters cannot be reliably determined except from J –V data measured in the dark.
      In addition, it must be verified that there are no nonohmic effects at any contacts or junc-
      tions, which cause the appearance of a second diode for which equation (13.7) does not
      account. Such nonohmic behavior is often observed at reduced temperatures [170, 172].
      Once it has been demonstrated that all these parasitic effects are negligible, or corrections
      have been made, then JO can be determined by a linear fit to a semilogarithmic plot of
      J + JL versus V –RS J and A can be determined from the slope of the derivative dV /dJ
      versus 1/J in forward bias [181], or both JO and A can be obtained by a least squares
      fit to equation (13.7). Finally, b can be determined from the temperature dependence of
      VOC as in Figure 13.18.
             It must be noted that most descriptions of transport and recombination ignore the
      effect of grain boundaries, implicitly assuming that grains are columnar and all transport
      can proceed without crossing grain boundaries. However, this is rarely, if ever, strictly
      true, so a comprehensive description of Cu(InGa)Se2 solar cells must account for the
      possibility of recombination at grain boundaries reducing current collection or voltage.
      The effect of grain boundaries can be expressed as an effective diffusion length, leading to
      the conclusion that grain-boundary recombination is small [27]. This can occur if the grain
      boundaries are doped more p-type than the bulk grains so that electrons are prevented
      from reaching and recombining at defects in the grain boundaries [169].
                                                                    DEVICE OPERATION            599

13.5.3 The Cu(InGa)Se2 /CdS Interface
     It may seem surprising that recombination at the Cu(InGa)Se2 /CdS interface does not
     limit VOC since, in processing Cu(InGa)Se2 solar cells, no special efforts are made to
     match lattices or reduce interface defects and the devices are typically exposed to air
     between the Cu(InGa)Se2 and CdS depositions. This can be explained by type inversion
     of the near-junction region of the Cu(InGa)Se2 induced by the band alignment and dop-
     ing [169, 182–184]. In this case, the Fermi level at the interface is close to the conduction
     band so that electrons in the near surface region of the Cu(InGa)Se2 are effectively major-
     ity carriers and there is an insufficient supply of holes available for recombination through
     the interface states. It has alternatively been proposed that doping due to Cd diffusion
     during the chemical bath deposition of CdS results in the formation of an n-type emitter
     and a p –n homojunction in the Cu(InGa)Se2 [136]. This would require the junction to
     remain very close to the Cu(InGa)Se2 /CdS interface to minimize recombination of carriers
     generated near the interface, and would therefore be very process-specific.
            The Cu(InGa)Se2 /CdS band diagram shown in Figure 13.19 demonstrates that the
     conduction-band offset EC between the CdS and the Cu(InGa)Se2 is critical for creat-
     ing the type inversion in the Cu(InGa)Se2 . In this diagram, the bulk Cu(InGa)Se2 layer is
     p-type with Eg depending on the relative Ga concentration, the CdS layer is n-type with
     Eg = 2.4 eV and is totally depleted, and the bulk ZnO n+ -layer has Eg = 3.2 eV. A thin
     HR ZnO layer between the n+ -ZnO layer and the CdS is also assumed to be depleted. Pos-
     itive EC indicates a spike in the conduction band, that is, the conduction-band minimum
     in the CdS is at higher energy than the conduction-band minimum of the Cu(InGa)Se2 .
     Figure 13.19 shows the case with EC = 0.3 eV and a −0.3 eV conduction-band offset
     between the ZnO and the CdS [52]. Models of current transport and recombination have
     considered the effect of EC [184–187]. These models show that if EC is greater than
     about 0.5 eV, collection of photogenerated electrons in the Cu(InGa)Se2 is impeded and

                           ZnO             CdS                 Cu(InGa)Se2
                        Eg = 3.2 eV     Eg = 2.4 eV            Eg = 1.2 eV


                                  ∆EC           JREC
                                                (p = n)

     Figure 13.19 Band diagram of a ZnO/CdS/Cu(InGa)Se2 device at 0 V in the dark. Note that the
     recombination current JREC is greatest where p = n in the space charge region of the Cu(InGa)Se2
     and not at the interface
600      Cu(InGa)Se2 SOLAR CELLS

      JSC or FF is reduced sharply. With a smaller spike, electrons can be transported across
      the interface assisted by thermionic emission [185]. On the other hand, for sufficiently
      negative EC the induced type inversion of the Cu(InGa)Se2 near the interface is elimi-
      nated and interface state recombination will limit VOC . An ODC layer at the surface of the
      absorber layer increases the band gap and primarily affects the valence band [188], so it
      may enhance type inversion near the junction. However, there is no convincing evidence
      that this layer exists in devices, so it is not shown in Figure 13.19.
             Owing to its importance in the electronic behavior of Cu(InGa)Se2 /CdS devices,
      several efforts have been made to calculate or measure EC with varying results.
      Band-structure calculations gave EC = 0.3 eV [189]. XPS and ultraviolet photoelectron
      spectroscopy (UPS) measurements of the valence band alignment indicate a positive EC
      between 0.2 and 0.7 eV [52, 190, 191]. These electron spectroscopy methods require
      ultrahigh vacuum conditions that necessitates that the CdS is deposited by vacuum
      evaporation. It is possible that the interface formation is different when CdS is grown
      by chemical bath deposition, for example, due to chemical interdiffusion, resulting in
      a different alignment of the conduction bands. Indirect measurements of the junction
      formed with chemical bath deposited CdS using a surface photovoltage technique
      gave EC = −0.1 eV [192]. Finally, inverse photoemission spectroscopy showed that
      substantial chemical intermixing occurs across the interface resulting in EC = 0 [193].

13.5.4 Wide and Graded Band Gap Devices
      While the highest efficiency devices generally have Ga/(In + Ga) ≈ 0.1–0.3 giving Eg ≈
      1.1–1.2 eV, significant effort has been made to develop high-efficiency solar cells based
      on wider band gap alloys. This is driven primarily by the expectation that wider band gap
      alloys will yield higher module efficiencies due to reduced losses related to the trade-off
      between higher voltage and lower current at maximum power. The resulting reduction in
      power loss, proportional to I 2 R, can be used to either (1) increase the module’s active
      area by reducing the spacing between interconnects or (2) decrease the optical absorption
      in the TCO layers since they can tolerate greater resistance. Wider band gap should give
      a lower coefficient of temperature for the device or module output power, which will
      improve performance at the elevated temperatures experienced in most real terrestrial
      applications. Wide band gap devices could also be used as the top cell in a tandem or
      multijunction cell structure.
              The wider band gap materials that have attracted the most attention for devices are
      Cu(InGa)Se2 and CuInS2 . CuGaSe2 has Eg = 1.68 eV, which is well suited for the wide
      band gap cell in tandem structures. CuInS2 has Eg = 1.53 eV, which could be nearly
      optimum for a single-junction device. The highest-efficiency devices based on CuInS2 are
      deposited with Cu-rich overall composition and then the excess Cu, in the form of a Cux S
      second phase, is etched away before CdS deposition [194]. Cu(InAl)Se2 solar cells have
      also been considered [195]. Since CuAlSe2 has Eg = 2.7 eV, the alloy requires smaller
      changes in relative alloy concentration and lattice parameter from CuInSe2 than the Ga
      alloys to achieve comparable band gap. The highest efficiency devices of different alloys
      are listed in Table 13.5.
             The effects of Ga incorporation on device behavior are not fully understood. The
      addition of a small amount of Ga to CuInSe2 increased the open-circuit voltage even when
                                                                                    DEVICE OPERATION        601

           Table 13.5 Highest-efficiency devices for different alloy absorber layers
           Material                  Eg    Efficiency           VOC           JSC        FF      Reference
                                    [eV]      [%]              [%]         [mA/cm2 ]    [%]
           CuInSe2                  1.02         15.4           515          41.2      72.6        [80]
           Cu(InGa)Se2              1.12         18.8           678          35.2      78.6         [1]
           CuGaSe2                  1.68          8.3           861          14.2      67.9       [196]
           CuInS2                   1.53         11.4           729          21.8      71.7       [197]
           Cu(InAl)Se2              1.16         16.9           621          36.0      75.5       [198]


                               14                                                        0.8





                                1.1        1.2          1.3          1.4      1.5      1.6

Figure 13.20 Efficiency ( ) and VOC (ž) as a function of Cu(InGa)Se2 band gap, varied by
increasing the relative Ga content, (From Shafarman W, Klenk R, McCandless B, Proc. 25th IEEE
Photovoltaic Specialist Conf., 763–768 (1996) [199]. The dashed line has slope VOC / Eg = 1

the Ga was confined to the back of the absorber and did not increase the band gap in the
space charge region [105]. The effect of increasing band gap in Cu(InGa)Se2 /CdS solar
cells on VOC and efficiency is shown in Figure 13.20. Efficiency is roughly independent of
band gap for Eg < 1.3 eV or Ga/(In + Ga) < 0.5 [165, 199]. With even wider band gap,
VOC increases to greater than 0.8 V, but the efficiency decreases. This indicates poorer
electronic properties of the Cu(InGa)Se2 absorber layer, which has two effects: voltage-
dependent current collection [165], which causes the fill factor to decrease, and increased
recombination [200], which reduces VOC below that expected from equation (13.9) [27].
The dashed line in Figure 13.20 shows a line with slope VOC / Eg = 1. Ideally, the
increase in VOC would have only a slightly smaller slope due to the dependence on
JL in the second term of equation (13.9). Admittance spectroscopy showed a correlation
between the recombination and the density of a defect with an activation energy ∼0.3 eV,
which increases with Eg [200]. Transient photocapacitance measurements showed a defect
band centered at 0.8 eV from the valence band, which moves closer to midgap for increas-
ing band gap and therefore becomes more efficient as a recombination trap [201]. As the
band gap becomes wider, type inversion of the absorber layer near the interface may no
longer occur and interface recombination can become more significant. Analysis of both
CuGaSe2 [202] and CuInS2 [203] solar cells showed that the low open-circuit voltages
602      Cu(InGa)Se2 SOLAR CELLS

      could be caused by either space charge or interface recombination, depending on the
      device preparation.
             Band gap gradients formed by controlled incorporation of Ga or S have been
      proposed as a means to increase device efficiency by separately reducing recombination
      and collection losses [19, 204–206]. A gradient in the conduction band from wide at
      the Cu(InGa)Se2 /Mo interface to narrow near the space charge region has been used to
      enhance minority-carrier collection [199, 206] and to reduce back surface recombination
      when the diffusion length is comparable to the film thickness [207]. Alternatively, a
      gradient from wide at the Cu(InGa)Se2 /CdS interface to narrow at the edge of the space
      charge region could reduce recombination and increase VOC . In this case, the smaller
      band gap in the bulk portion of the device can still enable high optical absorption and
      JSC [19, 206]. The most effective implementation of a surface band gap gradient may be
      the incorporation of S near the front surface [20] since the main effect is in lowering the
      valence band, instead of raising the conduction band as with Ga, and there should be less
      impact on collection of light-generated electrons.

      The competitiveness of a PV technology will primarily be governed by its performance,
      stability, and costs. The best Cu(InGa)Se2 cells and modules have demonstrated efficiency
      on a par with commercial crystalline silicon products. Long-term stability appears not to
      be a significant problem, as shown in field tests of prototype modules, but low-cost
      production remains to be demonstrated in practice.
              It is evident that thin films have the potential to be produced at very low costs.
      Moisture barriers of aluminum films that are deposited on plastic foils to be used, for
      example, in potato chip bags cost less than 0.01 $/m2 to produce. This particular example
      is at the low end of production costs and more advanced functional coatings are in general
      substantially more expensive to manufacture. Thin film coatings on architectural glass cost
      on the order of 1 $/m2 . Thus PV modules constructed from thin-film materials have the
      possibility for very low manufacturing costs. Whether Cu(InGa)Se2 module production
      will be able to achieve this low-cost potential will depend on how well the process
      technology fulfills the requirements for material costs, throughput, and yield.

13.6.1 Processes and Equipment
      Deposition processes can be either batch-type, in which a number of substrate plates are
      processed in parallel, or in-line, in which one substrate plate immediately follows the
      preceding one. In batch processing, a process step is completely finished before the next
      batch is started, whereas a substrate plate may enter an in-line process step before the
      previous substrate is finished in order to keep the line continuously running.
             One common view on volume production is that in-line continuous processing is a
      prerequisite for low costs. Fabrication of large-area thin-film products with physical vapor
      deposition is often made in a continuous or quasi-continuous in-line system. However, the
      cost of a batch process can be equally low, provided the throughput is large enough. For
      manufacturing of Cu(InGa)Se2 modules, this means that the CdS chemical bath deposition
                                                           MANUFACTURING ISSUES            603

can well fulfill low-cost production criteria even though it is normally a batch process.
Similarly, growth of the Cu(InGa)Se2 layer by batch selenization does not necessarily need
to be associated with higher costs than Cu(InGa)Se2 fabricated by in-line coevaporation,
provided the cycle time is short enough or batch size large enough.
       The commercial availability of large-area deposition systems depends on the spe-
cific process. Sputtering deposition is widely used for fabrication of large-area thin-film
coatings of various kinds, for example, in the glass industry. Similar processes are used
in the fabrication of most Cu(InGa)Se2 modules for the Mo back contact and the TCO
front contact, so the same type of equipment, available from a number of suppliers, can
be used.
       Sputtering is also typically used for deposition of the metal precursor films for fab-
rication of the Cu(InGa)Se2 layer by a two-step process. The selenization step, however,
requires specific custom-made process equipment. This could be selenization furnaces in
which batches of plates with the precursor layers are exposed to a selenium-containing
atmosphere or an in-line selenization chamber in which the plates are continuously
transported through an environment with elemental selenium and substrate temperature
control [208].
       Elemental coevaporation of the Cu(InGa)Se2 layer requires custom-made equip-
ment including specially designed evaporation sources for uniform deposition of large-area
substrates with accurate control. In-line evaporation using linear sources is a straightfor-
ward approach that is being developed at several laboratories and companies. An example
of such a piece of equipment is illustrated in Figure 13.21.
       The chemical bath deposition of CdS or Cd-free buffer layers is suitable for low-
cost batch processing, in that it is a surface-controlled process that requires a limited
solution volume. The equipment for dipping batches of Cu(InGa)Se2-coated substrate
plates is relatively simple and can be custom-made. The dry buffer deposition methods
under investigation have not been developed to a stage in which manufacturing is under
      Chemical vapor–deposited doped ZnO as an alternative to sputtering is typically
done as a batch process with a relatively small number of substrate plates deposited per

                   Raw                     Evaporation sources               Coated
                  plates                                                     plates

           Atmosphere      Vacuum                          Vacuum   Atmosphere

Figure 13.21 In-line coevaporation system for Cu(InGa)Se2 with linear evaporation sources above
the substrate plates and heaters below them [Courtesy of Zentrum f¨ r Sonnenenergie- und Wasser-
stoff-Forschung (ZSW)]. Reproduced by permission of Michael Powalla, ZSW Stuttgart, 2001
604      Cu(InGa)Se2 SOLAR CELLS

      run. Throughput will eventually become an issue. However, in-line CVD processes have
      been developed, for example, in the manufacture of amorphous silicon solar modules.

13.6.2 Module Fabrication
      Soda lime float glass is the substrate material that so far has given the best results in
      terms of both performance and reproducibility. It fulfills criteria on cost (3 $/m2 for
      4-mm-thick glass in large volumes), smoothness, and stability, so it is well suited for
      commercial production. One limitation that needs to be addressed in the development of
      production processes is that soda lime glass starts to soften above 500◦ C. At the same
      time, the best PV properties of Cu(InGa)Se2 are achieved at growth temperatures above
      500◦ C. Plastic deformation due to glass softening is not acceptable in a module-production
      process and careful optimization of the time–temperature profile is needed to minimize
      the deformation.
              Flexible substrate materials are attractive both for the possibility to make a
      lightweight flexible product with advantages for certain applications and for the possibility
      to deposit the thin-film materials in roll-to-roll processes, which are potentially cost-
      advantageous. Such roll-to-roll processing of semiconductor thin films was originally
      demonstrated with evaporation of CdS for solar cells [209]. The substrate materials
      that have shown promising results are polyimide, titanium, and steel [68, 69]. The
      drawbacks of polyimide are low-temperature tolerance, since the best polyimide films
      readily available can only withstand 400 to 450◦ C, and high thermal expansion. The main
      drawback of titanium and steel is their conductivity, which means that an electrically
      isolating layer is needed in order to allow monolithic series-interconnection of the cells.
      Such an isolation layer is not easy to make without local defects that will cause shunting
      of the cells. For these flexible substrate materials, sodium has to be supplied separately.
             An essential cost advantage with thin-film PV modules as compared to silicon
      wafer–based PV modules is the possibility of monolithic interconnection. This allows
      modules to be fabricated directly, instead of first making cells followed by tabbing and
      stringing to make the series interconnection as required for silicon-wafer solar cells. A
      typical monolithic interconnection is illustrated schematically in Figure 13.22. The most
      common way to make the patterning is by using laser ablation for the Mo patterning (P1)
      and mechanical scribing for the two subsequent patterning steps (P2 and P3).
             The final fabrication steps include attachment of electrical wires and buss bars.
      These are metal stripes that can be soldered, welded, or glued to contact areas near the
      edges of the substrate plates. Before lamination with a front cover glass, the thin-film
      layers are removed from the outer rim of the substrate plate in order to improve the
      adhesion to the lamination material, which is usually ethylene vinyl acetate (EVA). Edge
      sealing and framing finishes the product, but can be omitted for some applications.

13.6.3 Module Performance
      The evolution of record efficiencies as reported from the certified measurement labs
      is displayed in Figure 13.23. The module efficiencies lag behind the cell efficiencies
      but follow the same basic trend. There are additional losses associated with making
                                                                                       MANUFACTURING ISSUES                          605

                                               P1                                     P1


                                                           Glass substrate

                                                      P2                                       P2
                                                                                                                       HR ZnO

                                                             P3                                             P3


Figure 13.22 Schematic description of the manufacturing steps to make monolithic interconnec-
tions for thin-film Cu(InGa)Se2 PV modules

                            18   Cells > 1 cm2

                            17   Mini-modules < 100 cm2                                                                ÅSC
                            16   Modules > 3000 cm2
    Conversion efficiency

                            15                                                                                   ÅSC
                            14             Boeing NREL

                                                           Siemens          Solarex                         Siemens
                                                                                           Siemens      Siemens
                            10   ARCO
                                                              ÅSC - Ångstrom Solar Center, Uppsala University
                             9                                IPE - Institut fur Physikalische Elektronik, Stuttgart University
                                                              NREL - National Renewable Energy Laboratory, USA
                                 1991   1992        1993   1994      1995     1996         1997      1998    1999      2000   2001

Figure 13.23 Evolution of Cu(InGa)Se2 device record efficiencies in the past decade. All data
are taken from the Solar Cell Efficiency Tables periodically published in Progress in Photovoltaics
606       Cu(InGa)Se2 SOLAR CELLS

      series-interconnected modules instead of cells, both from series resistance and from
      inactive device area. In an optimized, conventional thin-film module design, these kinds
      of losses correspond to about 1% unit of efficiency. With a more advanced design using
      metal grids for interconnection, interconnect losses can be made nearly negligible [210].
      Another kind of difference between modules and record cells is associated with the free-
      dom to use higher process temperatures for cells that are not sensitive to deformed glass.
      These results are not necessarily relevant to module fabrication but indicate the potential
      of the materials.
             In a product the initial efficiency is of little interest if it deteriorates after some
      time in operation. Cu(InGa)Se2 modules fabricated by ARCO Solar and later Siemens
      Solar have shown stable performance in field tests over more than 12 years [3], as shown
      in Figure 13.24. On the other hand, severe degradation has been observed after exposure
      of cells to 85% relative humidity at 85◦ C for 1000 h [211], the so-called damp heat test,
      which is one of the certification tests in the IEC 61 646 protocol. While this test is rather
      severe and may not be relevant to thin-film modules, it shows the need for encapsulation
      techniques that minimize the exposure of the thin-film materials to moisture.
             The outdoor module performance demonstrated in Figure 13.24 shows that
      Cu(InGa)Se2 PV modules have the stability and performance to compete in any power
      application, be it stand-alone or grid-connected. Thin-film modules have a great advantage
      over silicon-wafer PV for consumer applications in which the power needed often is
      relatively small. The large substrate plates, which have a power of 40 WP or more, can
      easily be cut into smaller pieces, to essentially any power specification. This is much
      less costly than making small crystalline-silicon modules in which each cell has to be cut
      into pieces before assembling the modules. Additionally, the patterning structure of the
      interconnects can be designed to fit a large variety of shape and voltage requirements.
      Aesthetically, the solid black appearance of Cu(InGa)Se2 modules may be preferred to the
      nonuniform bluish appearance of the silicon-wafer modules in some building-integrated

                                          12    Modules tested outdoors at NREL
                                                      1988 module is 0.1 m2
                                                      others are 0.4 m2
                    Aperture efficiency




                                               1989              1992         1995     1998   2001

      Figure 13.24 Examples of outdoor testing results at NREL of Cu(InGa)Se2 modules showing
      stability over 12 years. Fluctuations in years 1992–1996 are due to changes in testing conditions.
      (Data courtesy of Shell Solar Industries)
                                                           MANUFACTURING ISSUES             607

             Finally, for space applications, Cu(InGa)Se2 thin-film solar cells offer potential
     advantages since the radiation tolerance is high as compared to crystalline-silicon solar
     cells [4, 5]. The potential to use a lightweight plastic substrate could lead to solar cells
     with very high specific power, that is, power divided by mass, which is critical for
     some space applications (see Chapter 10 for a more complete discussion). However,
     Cu(InGa)Se2 space solar cell technology has not yet reached a commercial stage.

13.6.4 Production Costs
     Material costs have direct and indirect components, and depend on the material yield
     of the deposition processes. The direct material costs, that is, the cost of the feedstock,
     will not be reduced by an increased volume of the production, depending only on the
     feedstock market price and how much material is needed in the film. The indirect costs,
     including preparation of sputtering targets or other source materials, will be reduced when
     production volumes are sufficiently large. The material yield, or fraction of the source
     material that ends up in the film, may be less than 50% for various thin-film processes.
     For sputtering, typically 30% of the target material ends up in the films.
            In addition to materials, the other main production cost for thin-film modules is
     the capital cost of the equipment. To first order, any large-scale automated deposition
     equipment will have comparable price. Therefore, the throughput or production capacity
     will be very important for determining the capital cost.
             Costs around 20 $/m2 for each thin-film deposition or process step may be accept-
     able in pilot production, but clear pathways toward costs in the range 1 to 5 $/m2 for
     large-volume production need to be identified. Throughput has a direct effect on cost.
     In an in-line process, this will depend on the substrate width and linear speed, which
     fundamentally depends on the deposition rate and desired thickness of the layer. If the
     deposition rate is relatively low, it can be compensated by having a long deposition zone
     in the system, for example, by having multiple targets in a sputtering system with only a
     relatively small increase in capital cost.
             All cost advantages for thin films are lost if the production is not completed with
     high yield. The overall manufacturing yield can be broken down into electrical yield
     and mechanical yield. The electrical yield reflects the module reproducibility since it is
     the fraction of the modules produced which fulfill minimum performance criteria. The
     mechanical yield is the fraction of the substrates entering the production line that make
     it to the end. Mechanical losses result from broken glass substrates or malfunctioning
     equipment. In general, the overall yield should be well over 80%.
           Another manufacturing cost is the energy usage. The energy payback time for
     Cu(InGa)Se2 modules is expected to be fairly low; four months has been estimated
     by Alsema and van Engelenburg [212], compared to three years for crystalline-silicon
     modules [213].
            Production-cost analyses result in a range of projected manufacturing costs. There
     are predictions of 1.5 to 2 $/WP for first-generation Cu(InGa)Se2 plants with a few MWP
     yearly capacity and projected costs of 0.4 to 0.6 $/Wp for large-volume manufacturing
     [214, 215].
608      Cu(InGa)Se2 SOLAR CELLS

13.6.5 Environmental Concerns
      One of the environmental issues related to the materials in Cu(InGa)Se2 modules is the
      availability of less common elements. The content of the critical materials in grams per
      kWP has been calculated assuming 12% module efficiency and the result is compared
      with the amount refined annually in Table 13.6 [216]. The fourth column expresses how
      much module power could be obtained from the amount refined annually and the last
      column shows a similar calculation based on the reserves of the various elements. Owing
      to uncertainties in estimates of reserves, or maximum resources, Table 13.6 just gives an
      indication of where, and at what level, potential problems in material supply may occur.
      It is clear that In is the potential bottleneck as regards primary material supply.
             CuInSe2 toxicity has been studied by administering it to rats [217]. Even at high
      doses negligible effects were detected. A lowest observed adverse effect level (LOAEL)
      of 8.3 µg/kg/day for humans was derived from these studies.
             The other substances that constitute Cu(InGa)Se2 modules are largely nontoxic
      except for Cd. Many aspects of its use in PV manufacturing have been studied by
      Fthenakis and Moskowitz [218]. Chemical bath deposition of CdS is the process step
      that presents the greatest health concerns due to the use of Cd, thiourea, and the genera-
      tion of waste solutions. In electrodeposition of CdTe, which also is a wet process using
      Cd precursors, it was found that the greatest health hazards from Cd are from dust gener-
      ated during feedstock preparation and from fine particles near the baths [218]. Biological
      monitoring at a process station showed that exposures can be maintained at a level that
      presents no risk to workers. Thiourea is a toxic and carcinogenic substance that also
      presents an exposure risk. Rinse water and dilute solutions of acids and Cd-compounds
      can be treated by a two-stage precipitation/ion exchange process. The Cd can be removed,
      and recycled, down to 1 to 10 ppb levels [218].
             Most Cu(InGa)Se2 processes use elemental Se, but the forms that are handled are
      solid shots or pellets that give off very little dust that could be inhaled. Elemental Se
      is considered to have a relatively low biological activity, but many compounds are very

                         Table 13.6 Critical materials in Cu(InGa)Se2 modules
                         with respect to primary supply (After Andersson B,
                         Azar C, Holmberg J, Karlsson S, Energy 5, 407–411
                         (1998) [216])
                         Element    Material   Amount     Amount     Reserves/
                                     content    refined     refined/    content
                                    [g/kWP ]   [kton/y]    content    [TWP ]
                                                          [GWP /y]
                           Mo        42          110        2600       130
                           Cu        17         9000      529 000     30 000
                           In        23         0.13         5.7        0.1
                           Ga         5         0.06         12         2.2
                           Se        43           2          46         1.9
                           Cd         1.6         20       12 500      330
                           Zn        37         7400      200 000      4100
                                                          THE Cu(InGa)Se2 OUTLOOK            609

    active and highly toxic. In particular, hydrogen selenide, a gas used in some selenization
    processes, is extremely toxic with an “immediately dangerous to life and health” (IDLH)
    value of only 2 ppm [218].
            There are also environmental concerns for the hazards during the operation of
    Cu(InGa)Se2 modules with one potential risk being the leaching of critical materials
    into rainwater. This only happens if a module is broken or crushed, so the normally
    well-encapsulated active layers are exposed. An experimental study of the emissions
    of toxic elements into rainwater from crushed CuInSe2 modules and into soil exposed
    to the water concluded that no acute danger to humans or the environment is likely to
    occur [219]. The main hazard during the active life of the CuInSe2 modules is related to
    fire accidents. A study of the potential risks associated with fires in PV power plants shows
    that they are very limited [220]. A fire in a commercial-size system could result in harmful
    concentrations up to 300 m downwind of the fire if most of the CuInSe2 materials are
    released. With release of 10% of the CuInSe2 materials, concentrations were not harmful
    even under worst-possible meteorological conditions. The study concluded that there are
    no immediate risks to the public from fires in sites with CuInSe2 modules.
            Concerns for disposal of Cu(InGa)Se2 have also been tested with respect to leach-
    ability. Zn, Mo, and Se are eluted in the highest amounts. On the basis of landfill criteria,
    CuInSe2 modules will pass requirements in both Germany and the United States [217].
    Because of the low volume and leaching rates of critical elements from CuInSe2 modules,
    they will not be classified as hazardous waste according to most US regulations [221].
           The evolution of environmental regulations, disposal options, and economics makes
    recycling increasingly important. In large-scale use of Cu(InGa)Se2 modules, the supply
    of rare elements, in particular indium, but also selenium and gallium, provides a further
    motivation for recycling. The cost of recycling may be favorably offset if module materials
    can be reclaimed. In particular, if the glass sheets can be salvaged and reused, there will be
    a net gain associated with the recycling procedure. Thus, recycling may be an important
    consideration in the choice of encapsulation method. Double glass structures are functional
    and may reduce the release of CuInSe2 materials during fires, but may increase the costs
    for recovering metals and reusing glass plates [221].

13.7 THE Cu(InGa)Se2 OUTLOOK
    Clearly, there has been tremendous progress in Cu(InGa)Se2 solar cells as evidenced by the
    high module and cell efficiencies fabricated by many groups, the range of deposition and
    device options that have been developed, and the growing base of science and engineering
    knowledge of these materials and processes. There is good reason to be optimistic that cell
    efficiencies greater than 20% will be achieved before long and that module performance
    and yield will continue to improve. Still, there is a lack of understanding of many of the
    critical problems associated with semiconductor processing and a need to devote time
    and research focus at both the laboratory scale, to address fundamental issues, and on the
    pilot line, to address equipment and scale-up problems and to validate processes.
           From their earliest development, CuInSe2 -based solar cells, along with other thin-
    film PV materials including Cu2 S, CdTe, and amorphous Si, attracted an interest because
    of their perceived potential to be manufactured at a lower cost than Si wafer-based PV.
610      Cu(InGa)Se2 SOLAR CELLS

      However, after more than 25 years of research and development of CuInSe2 , manufac-
      turing has only recently moved past the pilot-production stage and has not demonstrated
      any cost advantages. A fundamental question must be asked: what needs to be done
      to ensure that Cu(InGa)Se2 solar cell technology reaches its potential for large-scale
      power generation?
             Part of the answer is to address the critical need for the accelerated development of
      new manufacturing technology including improved deposition equipment and processes
      based on well-developed engineering models. Also, new diagnostic and process-control
      tools will have to be developed. This requires fundamental materials and device knowledge
      to determine what properties can be measured in a cell or module fabrication process that
      can act as reliable predictors of final performance. Better processes, equipment, and control
      based on a more solid knowledge base can directly translate to higher throughput, yield,
      and performance.
              There is also a critical need for continued improvement in the fundamental science
      of the materials and devices [222, 223]. Significant improvements in efficiency will only
      come from increased VOC so the chemical and electronic nature of the defects that limit
      it, and their origin, must be understood. This can contribute to a comprehensive model for
      the growth of Cu(InGa)Se2 , relating processing parameters to defect formation, junction
      formation, and device limitations. In addition, a fundamental understanding of the role of
      sodium and the nature of the grain boundaries and free surface needs to be developed. A
      greater understanding of the role of the CdS layer and the chemical bath process might
      enable alternative materials that do not contain cadmium and have wider band gap to be
      utilized with greater efficiency and reproducibility.
             A second fundamental question to be asked is: what might be the breakthroughs
      that could lead to the next generation of thin-film Cu(InGa)Se2 -based solar cells?
             Further development of wide band gap alloys to enable cells to be made with
      Eg ≥ 1.5 eV without any decrease in performance will have several benefits for module
      fabrication and performance as discussed in Section 13.5. In addition, development of a
      cell with Eg ≈ 1.7 eV is a prerequisite for tandem cells based on the polycrystalline thin
      films to be developed. A monolithic tandem cell has the potential to attain efficiencies
      of 25% or more. The CuInSe2 alloy system is ideally suited for such a structure since a
      CuInSe2 cell with Eg = 1.0 eV would make an ideal bottom cell with any of the alloys
      that increase band gap to 1.7 eV for the top cell. Even if a high-efficiency wide band
      gap cell is developed, such a structure will require the development of a transparent
      interconnect between the top and bottom cells and improvements in cell structure or low-
      temperature processes to allow the bottom cell to survive the subsequent processing of
      the top cell.
             Low-temperature processing of the Cu(InGa)Se2 layer without loss of efficiency
      in the final solar cell can have significant additional benefits. With lower substrate tem-
      perature, alternative substrate materials, like a flexible polymer web, can be utilized. In
      addition, lower TSS can reduce thermally induced stress on the substrate, allowing faster
      heat-up and cooldown, and decrease the heat load and stress on the entire deposition
      system. Similarly, there may be significant cost and processing advantages to a cell
      structure that enables the use of a Cu(InGa)Se2 layer much less than 1 µm.
                                                                           REFERENCES           611

           With all these challenges to improve the fundamental knowledge behind
   Cu(InGa)Se2 materials and devices and to develop new manufacturing technology and
   breakthrough advancements, research and development on Cu(InGa)Se2 and related
   materials remains exciting and promising. All of the reasons for the initial excitement over
   the potential for thin-film Cu(InGa)Se2 remain valid. The high efficiency, demonstrated
   stability, and tolerance to material and process variations give great hope that it will be
   a major contributor to our solar electric future.

     1.   Contreras M et al., Prog. Photovolt. 7, 311–316 (1999).
     2.   Tanaka et al., Proc. 17th Euro. Conf. Photovoltaic Solar Energy Conversion, 989–994 (2001).
     3.   Wieting R, AIP Conf. Proc. 462, 3–8 (1999).
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6 Chalcopyrite Based Solar Cells
Renier Klenk, Martha Ch. Lux-Steiner
Hahn-Meitner-Institut Berlin Gleinicker, Berlin


Chalcopyrite based solar modules uniquely combine advantages of thin film technology with
the efficiency and stability of conventional crystalline silicon cells . It is therefore believed that
chalcopyrite based modules can take up a large part of the photovoltaic (PV) market growth
once true mass production is started.
    The most important chalcopyrite compounds for photovoltaic applications are CuInSe2 ,
CuInS2 , and CuGaSe2 with bandgaps of 1.0, 1.5, and 1.7 eV, respectively. Together with related
materials they offer high optical absorption and a wide range of lattice constants and bandgaps
(Figure 6.1). The compounds can be alloyed to obtain intermediate bandgaps. Starting with
single crystals [2], chalcopyrite based solar cells have been under investigation since 1974. The
first chalcopyrite cells had a CuInSe2 absorber and therefore the technology is most advanced
for lower gap materials with a composition close to CuInSe2 . Today the efficiency of lab scale
thin film devices is close to 20 % [3], an efficiency comparable to the best multicrystalline
silicon cells. Many scaling up and manufacturing issues have been resolved. Pilot production
lines are operational and modules are commercially available. As of 2005 the market share of
chalcopyrite PV modules is not yet significant but major problems that might prevent further
commercialization have not been identified.

The attractive potential of chalcopyrite photovoltaic modules can be summarized by key points
which we will briefly illustrate here and assess in more detail in the reminder of the chapter:

r High efficiency.
r Stability.
r Low cost.
r Effective use of raw materials.

Thin Film Solar Cells Edited by J. Poortmans and V. Arkhipov
C 2006 John Wiley & Sons, Ltd


                Bandgap (eV)


                               1.5    CuInS2

                                     0.54      0.56          0.58          0.6         0.62
                                                      Lattice constant a

Figure 6.1 Bandgaps and lattice constants of selected chalcopyrites according to data compiled in [1]
(lines are a guide to the eye only).
r Short energy payback time.
r Adaptable to various applications.
r Large supporting research and development community.

   Chalcopyrites clearly offer the highest efficiency potential among all thin film technologies.
The record efficiency for a small, lab scale cell is close to 20 %, using just a single layer
antireflective coating and a standard metal grid but none of the complex concepts that have
been used to produce record silicon cells. Submodule efficiencies are at almost 17 % [4], and
square foot and larger modules range from 14 to 12 % efficiency [5–7]. It is conceivable that
the maximum efficiency can be increased further. A significant boost, exceeding even the
theoretical limit for silicon, can be expected from the development of multijunction cells . The
excellent performance is notable not only under standard reporting conditions but also when
assessing monitoring data from outdoor installations.
   In contrast to amorphous silicon based cells chalcopyrite devices do not show any degra-
dation under illumination. Outdoor testing indicates that achievable product lifetimes may be
comparable to those of conventional photovoltaic modules.
   The low cost potential is roughly comparable to that of other thin film technologies and is
rooted in the use of inexpensive substrates, effective use of raw materials, high throughput,
and large area deposition at low temperatures as well as monolithic interconnection.
   Apart from the substrate, the total thickness of a chalcopyrite cell, including all films, is in
the range of 2 to 4 μm, which implies that the raw material usage is only a tiny fraction of the
material input for a silicon cell. Mass production will not be limited by the availability of raw
   The energy payback time (EPBT) is obviously an important parameter when considering
how far photovoltaics can contribute to the future energy supply. The much lower thermal
budget of thin film preparation (lower process temperatures as well as short process times)
leads to a significant benefit.
                                            CHALCOPYRITE BASED SOLAR CELLS                        239

    As we will show in this chapter, there is considerable flexibility concerning the choice of
components of a chalcopyrite cell or module as well as concerning the preparation methods for
these components. It is therefore possible to design products with an optimum efficiency/cost
trade-off for various applications, with power demands ranging from mW to MW and with
illumination intensity ranging from indoor, low level to high level under concentration. Chal-
copyrite cells can be grown on rigid as well as flexible substrates. They perform well in chal-
lenging environments because they are mechanically robust, can operate in a wide temperature
range, and can tolerate high radiation levels.
    The chalcopyrite technology is supported by a networked international research commu-
nity comprising universities, research institutes, and companies with experience and long term
commitment. This ought to guarantee that any problems surfacing in industrial implementation
can be attacked with the necessary background knowledge and that new developments are in
the pipeline for future products. The large financial risk of implementing mass production
facilities is limiting the production volume. However, the growing involvement of compa-
nies developing deposition equipment and fabrication infrastructure provides a solid basis for
ongoing commercialization.

The solar cells consist of a number of films which are deposited onto a rigid or flexible substrate
(Figure 6.2). The first film, typically molybdenum, serves as a nontransparent back-contact.
It is covered by the actual chalcopyrite film. This p-type film absorbs most of the light and
generates the photocurrent (absorber). The heterojunction is formed by depositing a very thin
n-type buffer layer (typically CdS) and an n-type wide gap transparent front contact (usually
heavily doped ZnO).

Figure 6.2 Scanning electron micrograph of the cross section of a typical chalcopyrite solar cell with
Cu(In,Ga)Se2 (CIGSe) absorber (substrate not shown).
240            THIN FILM SOLAR CELLS

                                                                                      500 °C


      Float glass
                         Mo                            Cu             In
                                Strukturierung 1
                Mo Sputtering   Laser Scribing     Precursor Sputtering          Reactive RTP

           Chemisches Bad   Strukturierung 2                                          Verkapselung
           Chemical       Mechanical Scribing        TCO        Mechanical Scribing     Encapsulation
              Bath            (Pattern 2)          Sputtering       (Pattern 3)

Figure 6.3 Typical sequence of processes to prepare a chalcopyrite photovoltaic module [8]. The
sequence shown here is based on the sequential approach (see text) using rapid thermal processing
(RTP). Reprinted from Thin Solid Films, 481–482, R. Klenk, J. Klaer, R. Scheer, M. Ch. Lux-Steiner,
I. Luck, N. Meyer, U. R¨ hle, 509, Copyright (2005) with permission from Elsevier.

    Thin films with properties suitable for photovoltaic applications can be prepared by a variety
of processes. In pilot lines the absorber is grown by multisource evaporation or by a sequential
approach (reactive annealing of metal films). To provide just one example, Figure 6.3 shows a
sequence of processes necessary to produce a CuInS2 based module using the latter method. In
the following paragraphs we will, for each of the films, introduce the most common state-of-the-
art methods as they are the basis of today’s production lines. Methods still under development
and aiming at future significant cost reduction are described in Section 6.6.

6.3.1 Absorber
The most advanced absorber materials are based on CuInSe2 (CISe). Due to its rather
low bandgap (1.0 eV) a common practice has evolved where gallium is added to obtain
Cu(In,Ga)Se2 with a wider bandgap of approximately 1.15 eV. Due to reasons outlined be-
low, sulfur is incorporated additionally in sequential processes to obtain Cu(In,Ga)(Se,S)2
(CIGSSe). All these absorbers require an optimized sodium concentration for optimum film
properties (see Section Another line of development starts from CuInS2 (CIS). In
contrast to the other compounds, this base material is grown under copper excess and it also
has a higher bandgap (E g = 1.5 eV). Addition of further elements such as sodium or gallium
is therefore not strictly necessary. The most commonly used absorber preparation methods
are multisource evaporation and sequential processes. Evaporation offers a much more direct
control of film formation, including deliberate depth profiles for bandgap engineering as well
as making optimum use of the fluxing activity of Cu(S,Se) phases. This is probably the reason
                                                            CHALCOPYRITE BASED SOLAR CELLS         241

why evaporation still yields the highest cell performance. In particular, unlike evaporation, the
sequential processes lead to an aggregation of gallium close to the back contact which causes
low gallium content in the active cell region and, consequently, a bandgap below the optimum
value. The addition of H2 S to the annealing environment has helped to overcome this problem
to a certain degree but it also adds to the overall complexity of the processes and the device.
In evaporated films the standard bandgap (in the order of 1.15 eV) can be achieved merely by
incorporating sufficient amounts of gallium. Multisource evaporation

Thin films can be grown in a straightforward manner by coevaporating the constituent elements
onto a heated substrate. The stoichiometry (concentration of VI element relative to the metals)
is handled by a group VI overpressure which has to be maintained in the initial stage of cooling
down the substrate. On the other hand, the molecularity (ratio of group I metal versus group
III metal concentration) has to be adjusted by tight control of the metal source temperatures.
Single crystal substrates with a suitable lattice constant and surface termination can be used
and will result in epitaxial growth of the thin film.
    The morphology and other properties of the resulting film depend strongly on the molecu-
larity. Copper rich films exhibit larger grains. They are a mixture of chalcopyrite with a close
to ideal composition and Cu-VI binary phases, typically found at the surface after cooling
down the sample. The chalcopyrite grown under these conditions is characterized by lower
defect density and reduced compensational doping in comparison to material grown without
Cu excess (Figure 6.4). These observations suggest that the binary phases play an active role in
the growth mechanism, also with regard to incorporation of the VI element [9]. Films with an
overall molecularity close to unity are often found to be inhomogeneous on a scale of several
μm due to localized segregation of Cu-VI binary phases.
                 Conductivity (Ω cm)-1



                                           0.48   0.49        0.50    0.51       0.52   0.53
                                                         Ga/(Cu+Ga) before etching

Figure 6.4 Lateral conductivity of CuGaSe2 thin films. Films grown under Cu rich conditions show
higher effective doping than those grown under Cu poor conditions. The Cux Se segregations were removed
by chemical etching prior to the measurement.

                                               Line Sources
                                       Se    Cu Se In Ga Se

           Load Lock                                                                   Load Lock

                  Substrate Carrier         Heaters

Figure 6.5 Schematic view of an in-line evaporation system. Substrates are supported by a carrier and
transferred into the evaporation chamber via a load lock. The substrate temperature is raised by an array of
heaters. Before reaching the maximum temperature they pass the copper source and then, after reaching
the maximum temperature, the indium and gallium sources. The substrate carriers are gradually cooled
down prior to leaving the evaporation chamber via the load lock at the right hand side.

    Notwithstanding the small grain size, efficiencies in the order of 14 % are readily achieved
with slightly Cu poor Cu(In,Ga)Se2 films prepared by this simple coevaporation approach. On
the other hand, several schemes have been developed to exploit the growth assistance that goes
along with excess copper. In the bilayer approach [10], a Cu rich coarse grained seed film is
grown first. The Cu rate is then diminished so that the Cu excess present in the seed films is
gradually consumed in this second stage, ending up with a slightly Cu deficient film. The idea
here is to combine the superior properties of Cu rich films with the absence of second phases
in Cu poor films. It is clear, however, that the growth mechanism cannot be sustained once the
Cu VI phases have been consumed which may lead to disrupted growth and new nucleation.
The advantage of this approach lies in the fact that it can be translated into an industrial
inline process where the moving substrate passes first the Cu, then the In and Ga sources
(Figure 6.5). Other schemes start from Cu poor films or even films without any copper and add
more copper in a second stage (inverted bilayer). If enough copper is delivered in this stage
the film becomes Cu rich and recrystallizes. In this case a third stage is needed to again reduce
the Cu content and achieve single phase material. This scheme is generally known as the three
stage process [11] and has resulted in the highest efficiency of lab scale devices so far (close to
20 %). In principle, there is again the problem of maintaining the growth mechanism after the
excess copper has been consumed in the third stage. However, the three stage process enables
precise control. The transition points where the film enters and leaves the Cu rich regime are
observable through changes in substrate temperature [12] (or power delivered to the heater to
maintain a constant temperature) which is often used for in situ process control. Monitoring
the intensity of light reflected off axis from the substrate is an alternative method for process
control (laser light scattering [13]). The three stage process results in a certain depth profile
of the Ga/(Ga+In) ratio which is believed to contribute to the excellent performance of cells
prepared from these absorbers (bandgap engineering). Sequential processes

Sequential (two step) processes have been developed as an alternative approach to absorber
formation [14, 15]. Here, a metallic precursor is typically deposited by sputtering. Sputtering
                                          CHALCOPYRITE BASED SOLAR CELLS                      243

of pure gallium is problematic due to its low melting point. Copper/gallium alloyed targets are
typically used instead. The chalcopyrite is formed in the second step by exposing the precursor
to a chalcogen containing atmosphere at elevated temperatures (selenisation/sulfurisation). The
method is particularly attractive for production. A process well established in industry is DC
sputtering and off-the-shelf equipment is readily available. It is characterized by good repro-
ducibility and large area uniformity of the thicknesses of the individual films. Consequently,
the important Cu/III ratio can be tightly controlled in this first step. The high temperature,
corrosive environment which is potentially problematic with respect to equipment degradation
over time is limited to the second step. Here, it is less critical because this second step mainly
affects the chalcogen stoichiometry which is to a great extent self adjusting.
    The second step can be carried out in a tube furnace or by rapid thermal processing (RTP).
Annealing in a tube furnace (using mixtures of an inert carrier gas and H2 S and/or H2 Se reac-
tive components) is typically a slow batch type process where several substrates are processed
simultaneously and where the substrate size is somewhat restricted by the maximum available
diameter of the furnace tubes. More recently rapid thermal processing furnaces have been
introduced [16, 17]. In one of the processes selenium is evaporated onto the metal precursor
(stacked elemental layers) before the high temperature annealing whereas sulfur is still intro-
duced as H2 S during the RTP. Chalcopyrite formation in the two step process depends largely
on the thermodynamics and phase formation kinetics of the material system. These funda-
mental properties have therefore been investigated in detail in order to be able to optimize the
processes [18, 19].
    In addition to producing low gap Cu(In,Ga)(S,Se)2 absorbers, two step processing has also
been found particularly suitable for preparing CuInS2 (E g = 1.5 eV) [17]. In this case it is
possible to simply place pieces or powder of sulfur next to the substrate in the RTP furnace,
thereby eliminating the need for chalcogen evaporation in a separate process as well as the use
of any toxic gas. Very short annealing periods (a few minutes at top temperature) are achievable
due to the growth assistance by Cux S using Cu rich (Cu/In = 1.2–1.8) precursors. Films grown
under these conditions exhibit high p-type conductivity, therefore the sodium concentration
(see Section is not critical [20]. Blocking layers and sodium precursor films are not
used. As already mentioned, it is also not necessary to incorporate additional elements for
bandgap adjustment. Sodium

Recognizing the important influence of sodium on device performance had been a major
breakthrough in the development. Sodium appears to influence the growth mechanism leading
to superior morphology as well as higher effective p-type doping. In terms of cell perfor-
mance the latter seems to be the more important aspect. The observed increase in open circuit
voltage is quantitatively consistent with the measured increase in net doping in good cells
where recombination in the space charge region is dominant [21]. The fact that an increase
in device performance is also achieved by diffusing sodium into an already prepared film is
also a strong indication that the electronic effects are more significant than the morphology
   Especially in the slower evaporation processes, sodium diffusing from a soda lime glass
substrate through the back contact can yield the required sodium concentration in the absorber
film. If the substrate does not contain sodium, sodium salts (e.g., NaF) can be coevaporated

or deposited as a precursor film onto the back contact before depositing the absorber. The
latter approach, in combination with a diffusion blocking layer underneath the back contact is
sometimes also used with soda lime substrates to achieve a more accurate control of sodium
concentration, especially in fast sequential processes.
    Films prepared with excess copper are characterized by large grains as well as high con-
ductivity and, hence, do not require sodium doping. At the moment, this is technically relevant
only for CuInS2 based modules.

6.3.2 Contacts Diffusion barrier and back contact

Sputtered molybdenum is the most frequently used back contact. The inherent stress in the film
can be adjusted over a wide range through the pressure of the working gas in the sputter process
[23]. Optimized films adhere very well to glass or other substrates and laser or photolithographic
patterning is straightforward. A possible alternative to sputtering is e-gun evaporation which
so far has been used only in laboratory scale preparation but may also have cost advantages in
large scale production.
   The actual contact to the chalcopyrite is complex and may involve a Mo–chalcogenide
intermediate layer which forms during absorber preparation. At low temperatures the I (V )
curve often deviates from that of an ideal diode under forward bias. This bend-over of the
curve has been attributed to blocking at the back contact of the cell. However, no voltage drop
across the back contact could be found with appropriate test structures [24]. There are other
models to explain the I (V ) curve bend-over without involvement of a blocking contact [25].
Alternative contact materials for improved optical reflection and novel device configurations
are under investigation (Section 6.6).
   As already mentioned, diffusion barriers (silicon oxide, silicon nitride) deposited onto
the glass substrate before the molybdenum are not strictly necessary but can be used for a
more precise control of sodium doping. On the other hand, metal foil substrates often require
additional coatings underneath the molybdenum for blocking of impurities as well as substrate
planarization and isolation [26].

          Load Lock     Sputter Chamber      Se Evaporation          RTP Furnace        Load Lock

                      CuGa        In            Substrate            Heaters

Figure 6.6 Schematic view of an in-line sequential system. Copper and gallium are sputtered from an
alloy target. After sputtering indium, the substrate is transferred to an evaporation chamber and coated
with selenium. The completed stack is annealed in the rapid thermal processing furnace. No substrate
heating is required during metal and selenium deposition.
                                           CHALCOPYRITE BASED SOLAR CELLS                        245 Buffer

The thin (typically 50 nm) buffer layer is grown from a chemical bath [27]. Typical solutions
contain a Cd salt, thiourea as a sulfur source and ammonia in an aqueous solution. The substrate
is immersed in the cold solution. The solution is then heated to 60–80 ◦ C. The thiourea hy-
drolyzes and cadmium and sulfur ions recombine to form CdS. The films grow either directly
at the substrate or nanoparticles are formed in the solution and deposited onto the substrate.
Depending on the deposition conditions, and due to the aqueous environment, the film may
contain significant amounts of oxygen and hydrogen. Chemical bath deposition (CBD) of CdS
is very reproducible and yields good cell performance on any chalcopyrite absorber. Window

The preferred window or TCO (transparent conductive oxide) film consists of ZnO deposited
by sputtering or metal organic chemical vapor deposition (MOCVD). This film needs to have
a high lateral conductivity in modules to avoid ohmic losses. It is therefore highly doped with
aluminium or gallium (sputtered films) or boron (MOCVD). However, depositing a film with
low lateral resistance directly onto the buffer increases the negative influence of local defects
(such as pin holes [28]) and local fluctuation of absorber properties (e.g. the bandgap [29]).
This can be avoided by first depositing a thin (in the order of 100 nm) ZnO film with lower
conductivity, i.e., by sputtering from an undoped target or by adding oxygen to the working
    The window layer contributes significantly to the module cost. Low resistivity is therefore
desirable to minimize the film thickness. In practice, the resistivity of large area ZnO thin films
is in the order of 5 × 10−4 cm and cannot be significantly improved because higher doping
reduces the electron mobility and causes poor transmission due to free carrier absorption. It has
been argued that the film properties are very close to physical limits [30] and that, in the long
run, ZnO could be replaced by other materials with higher mobility (lower effective electron
mass). Approaches to reduce the cost of ZnO preparation include new methods to fabricate
ceramic targets or to use reactive sputtering from metallic targets.
    The ZnO film plays an important role for module stability in accelerated lifetime testing
under damp heat conditions which forms a part of the EN/IEC 61646 certification. The lateral
resistance tends to increase, giving rise to fill factor losses. It is therefore mandatory to optimize
ZnO preparation not only with respect to the as-grown properties but also by taking into account
the degradation in damp heat. Monolithic integration and encapsulation

Manufacturing of modules adds some process steps to the cell preparation outlined above.
The module is divided into cells which are connected in series by monolithic integration. The
connection is made from the molybdenum back contact to the TCO during TCO deposition
(Figure 6.7). A front metallization has been suggested but is normally not applied. The common
scheme requires three patterning steps: an isolation scribe in the molybdenum (P1), scribing the
absorber to create a gap which is later filled by TCO (P2), and an isolation scribe of the complete
cell structure down to the molybdenum (P3). While the preferred tool for P1 patterning is a

Figure 6.7 Optical microscopy image of a laser scribed line in a molybdenum film on glass. Pulse
frequency of the laser and scribing velocity have been adjusted for optimum overlap of pulses.

pulsed Nd-YAG Laser (Figure 6.8), photolithographic patterning is also possible. After laser
patterning the substrate is again subject to wet cleaning with rotating brushes to remove loose
particles. P2 and P3 patterning are carried out by mechanical scribing. P2 patterning can be
carried out before or after deposition of the undoped ZnO layer, the latter method may give a
better contact.
    The interconnection area constitutes a loss in active area and hence photocurrent. In princi-
ple, the scribe lines themselves should be as narrow as possible, the distance between P1, P2 and
P3, respectively, should be minimal, whereas the distance between two interconnection areas
should be maximal. In practice, a certain width of scribe line has to be maintained for sufficient
isolation (P1, P3) and contact resistance (P2). Hence the total width of the interconnection is
in the range of 0.5 to 1 mm. In addition, the allowable distance between interconnection areas
is limited by the lateral resistance of the ZnO film. The distance can be increased when using
a thicker ZnO film. However, a thicker film, apart from being costly, also causes photocurrent
losses due do its reduced transparency. Consequently the typical distance is in the range of 5 to
10 mm which results in an area loss due to interconnections of approximately 10 %. Wide gap
absorbers offer more flexibility in designing the module due to their reduced current density.
Typical photocurrent densities under full illumination are 42 mA/cm2 at a bandgap of 1 eV

                                P1          P2            P3



Figure 6.8 Schematic cross section of the cell interconnect in monolithic integration. This figure shows
the variant where the P2 scribing is carried out after the deposition of the undoped ZnO.
                                                                                 CHALCOPYRITE BASED SOLAR CELLS   247

                 Integral fluorescence intensity [a.u.]

                                                                after DH

                                                                                   before DH
                                                                           P2                        P3

                                                          1.6       1.4    1.2      1.0        0.8    0.6   0.4
                                                                           Lateral position [mm]

Figure 6.9 Locally resolved sulfur fluorescence (XES) intensity across an interconnect test structure
before and after accelerated ageing in damp heat (DH). The position of the scribe lines is indicated
(P2, P3) [31].

(CISe) and 22 mA/cm2 at 1.5 eV (CIS), respectively. A computer simulation has been made
available to optimize the module patterning for a given set of film properties [32].
    The interconnection appears to play a certain role in module degradation in accelerated
aging tests (damp heat). Locally resolved X-ray emission spectroscopy (XES) scans (Figure
6.9) on specially prepared test structures (larger scribe line distance, reduced ZnO thickness)
results in a preliminary model of contact degradation. Before damp heat, a large sulfur signal
is observed within the P3 scribe line which is due to sulfurisation of the molybdenum surface
(which occurs during absorber formation). A smaller signal is observed within the P2 scribe
due to the signal being attenuated by the ZnO film. This latter signal increases significantly
upon damp heat treatment. The spectra suggest that this is due to the formation of ZnSO4 ,
i.e., damp heat causes a chemical reaction between MoS2 and ZnO thereby deteriorating the
contact in the P2 scribe line.
    Concerning the described standard procedures there is still room for improvement because
patterning unfortunately interrupts the in-line vacuum processing. Patterning is also critical
with respect to the throughput (cycle time).The final glass–glass laminate is produced using
EVA foil and standard laminators as in the silicon technology. Encapsulation is critical for
passing the accelerated aging tests.

In terms of modeling, the chalcopyrite based solar cell is a quite complex device compris-
ing a number of polycrystalline compound semiconductor films and several heterointerfaces.
Nevertheless, tremendous progress has been achieved in measuring, modeling and understand-
ing various aspects of the device. Novel characterization methods specifically adapted to the
problems at hand have been developed and introduced. Examples are Kelvin probe force mi-
croscopy (KPFM) to measure work functions with submicron lateral resolution (Figure 6.10),

Figure 6.10 KPFM measurement [34] of a mechanically polished cross section of a ZnO:Ga/
(Zn,Mg)O/CIGSSe thin-film heterostructure: a) topography, b) work function measured simultaneously
with topography. Bright colour corresponds to high work function. c) is a plot of height, work function,
and electrical field along the path indicated in a). Reprinted from Thin Solid Films, 481–482, Th. Glatzel,
H. Steigert, S. Sadewasser, R. Klenk, M.Ch. Lux-Steiner, 177 , Copyright (2005) with permission from

inverse photoemission spectroscopy (IPES) to directly measure conduction band line ups, and
X-ray emission spectroscopy (XES) for the analysis of buried interfaces. In general terms,
quantitatively extracting material and device parameters from a measurement result requires
a model that describes the correlation between the property assessed by the measurement and
the underlying physical parameters. Due to the complexity of compound polycrystalline semi-
conductor films, the models are often an approximation. The extracted parameters have to be
interpreted as effective parameters bound to the specific model. They can be useful for com-
paring different samples quantitatively but can be misleading when used for calculations in a
different context. Numerical modeling has emerged as a useful tool for a better understanding
of the device. It is important because analytical approximations that can be safely made for
other solar cells types are not valid in the chalcopyrite cell. In conclusion, the materials science
of chalcopyrites, and solar cells based on them, is a wide field under active development. Here,
we have to limit the discussion to a few selected topics and the reader is referred to the literature
for more in depth information (the overview given in [33] is a good starting point).

6.4.1 Cell concept

The high optical absorption of the direct semiconductor chalcopyrites makes very thin ab-
sorbers feasible. However, it also means that the incident sunlight is absorbed close to the
surface. Assuming it would be possible to dope chalcopyrites in a well controlled manner it
would still be challenging to reach high efficiency with a homojunction solar cell. Depending
on surface passivation, the major part of carriers generated between the surface and the pn
junction would be lost due to surface recombination. This problem is avoided by introducing
the window/absorber heterojunction concept. Due to the wide bandgap of the window, the
absorption is shifted away from the surface to the internal interface. Even assuming no surface
                                          CHALCOPYRITE BASED SOLAR CELLS                      249

passivation and a very high surface recombination velocity, the losses are nevertheless small
because only an insignificant part of the light is absorbed in the window.
   On the other hand, the internal heterointerface might itself cause recombination, leading
not only to photocurrent loss but also to high bucking currents and consequently low open
circuit voltage. But calculations show that interface recombination does not necessarily have a
significant impact on cell performance [35]. The severity of interface recombination depends

r The density of recombination centers at the interface.
r The doping of absorber and window.
r The type and density of charged interface states.
r The conduction band line-up.

    A low density of interface states is always advantageous but may in practice be difficult to
achieve because the large area technology cannot be compared to the ultra clean environment
required for defect free epitaxial growth. The other, more feasible approach to lowering re-
combination lies in minimizing the density of either electrons or holes at the interface, which
requires appropriate doping, band line-up and interface charge. In terms of bucking current
(open circuit voltage) minimizing the density of either type of carrier yields comparable results.
Considering the photogenerated carriers it is, however, mandatory that the electrons collected
from the absorber are majority carriers at the interface (inverted interface). In conclusion, the
structure should be an n+ window/p absorber heterojunction where the Fermi level (E F ) at the
interface is close to the conduction band and where the Fermi level intersects midgap energy
at a small distance from the interface in the absorber. The interface charge should be positive
(donor like defects) to assist in establishing this structure. Likewise, the energetic position
of the absorber conduction band edge (E C ) should be slightly lower than that of the window
(spike) because the opposite situation (cliff) tends to push the absorber conduction band away
from the Fermi level. Furthermore, in the cliff type band line-up electrons from the window side
of the interface could recombine with holes from the absorber side which leads to a reduced
    In the actual chalcopyrite cell these design considerations must be implemented by the buffer
layer and its preparation technique. Measurements of the band line-up and Fermi level position
at the interface (see [37] for a review of chalcopyrite surface and interface properties) are not
straightforward and the results show significant variations [38–41]. The band diagram shown
in Figure 6.11 is believed to be a reasonable approximation and shows the type inversion of the
interface as required according to the above considerations. There are theories and supporting
measurements that indicate an inherent widening of the absorber bandgap towards the surface
[39, 42, 43]. It has been speculated that this effect is due to the segregation of distinct Cu
poor ordered vacancy compound (OVC) but there is no conclusive evidence concerning the
actual structure. If this widening is mainly due to a shift of the valence band edge (as shown
in Figure 6.11) this also decreases the hole density at the interface and lowers the interface
recombination even further. Obviously, this is only helpful for device performance when the
transition from bulk chalcopyrite to surface bandgap widening occurs without introducing a
high defect density. However, lattice matching between the bulk and surface phases may be
absent in wide gap chalcopyrites [44].

                                           ZnO:Al ZnO CdS Cu(In,Ga)Se2 EG = 1.2 eV

                     Energy (eV)        EC

                                   -1                      OVC


                                   -3   EV

                                    0.25         0.5        0.75        1       1.25
                                             Distance from front contact (μm)

Figure 6.11 Tentative calculated [36] band diagram of a ZnO/CdS/Cu(In,Ga)Se2 heterojunction as-
suming a widening of the bandgap at the absorber surface due to an ordered vacancy compound (OVC).

6.4.2 Carrier density and transport

Without deliberate doping, the majority carrier (hole) density in chalcopyrite thin films is
usually in the range of 1014 −1017 cm−3 and well suited for the application. Oxygen and
sodium are the most common impurities known to influence the carrier density. The density
of donors is found to be almost comparable to the acceptor density (compensation, self-
compensation [45]). The effective mobility in polycrystalline films at room temperature is in
the range 1–100 cm2 /Vs. A complete set of temperature dependent conductivity and Hall effect
data including polycrystalline as well as single crystal samples is only available for a certain
(prepared with Cu excess) type of CuGaSe2 . They are summarized and discussed in [46].
   Charge trapped at grain boundaries gives rise to space charge regions extending into the
grains. The resulting band bending leads to potential barriers. Hall effect measurements suggest
barriers in the order of 50 to 150 meV which can be explained by a charge density in the order
of 1012 cm−2 at the CGSe grain boundary [47]. Kelvin probe force microscopy in ultra high
vacuum is a relatively new characterization tool well suited to further clarify the grain boundary
models. The high lateral resolution of the microscope allows work function measurements
across individual grain boundaries. An average drop in the work function of 110 meV across
grain boundaries has been measured in good agreement with the Hall data [48]. The electrical
fields in the vicinity of grain boundaries would sweep minority carriers into the grain boundary.
However, excessive grain boundary recombination is in contradiction to the high photocurrents
which are observed even in fine grained films. Alternative models of the grain boundary have
been proposed, based on band structure calculations [49]. A comparison of the impact of either
type of grain boundary upon device performance has been performed by two-dimensional
numerical calculations [50]. Kelvin probe force microscopy results on the differences in the
local variation of the work function with illumination (surface photovoltage) suggest that both
aspects, i.e., trapped charge as well as the band structure, play a role in defining the grain
boundary properties [51].
                                           CHALCOPYRITE BASED SOLAR CELLS                       251

   There are a number of studies concerning time constants and pathways for radiative re-
combination as deduced from steady state photoluminescence and photoluminescence decay.
Radiative recombination appears to be dominant in good quality polycrystalline films at low
(8.5 K) temperature [52]. Free carrier lifetimes in the order of some nanoseconds have been
measured. Defect related lifetimes were significantly higher, presumably due to trapping of
carriers. Recent photoluminescence investigations [53] reveal a shallow donor (D) and two
shallow acceptor levels (A1, A2) which are present throughout the Cu(In, Ga)Se2 compounds.
The levels are found at slightly increasing depth when going from CuInSe2 (D = 10 meV, A1 =
40 meV, A2 = 60 meV) to CuGaSe2 . (D = 13 meV, A1 = 60 meV, A2 = 100 meV). The relative
concentration of the acceptor level is correlated with Cu/(In+Ga) ratio used during preparation
of the film. For CuGaSe2 , the same set of shallow defects can explain the photoluminescence
as well as the Hall measurement data [54]. It may be tempting to assign these levels to the
point defects of a ternary system whose formation energies and depth has been theoretically
calculated [55] , but the assignment remains uncertain in view of more complex defects and
impurities which might also play a dominant role.
   At room temperature the photoluminescence decays with a time constant of several tens
of nanoseconds due to nonradiative recombination. The majority of measurements described
in the literature and concerning minority carrier transport is assessing the minority carrier
diffusion length. The latter can be extracted from electron beam induced current (EBIC) [56]
and quantum efficiency measurements [57]. Typical values are in the range of 1–2 μm for
good quality films. This implies that the diffusion length and the extension of the space charge
region are comparable, which can cause problems in extracting both parameters from a single
measurement. The ambiguity can sometimes be resolved by measuring at varied applied bias
voltages and using an analytical approximation to calculate the field zone as a function of bias

6.4.3 Loss mechanisms
Depending on absorber material and bandgap, record efficiencies for chalcopyrite based so-
lar cells are in the range of 10 to almost 20 %. Optical [59] and contact related losses are
small, at least in small area cells with front contact grids. They need to be considered for
modules (without grids [4, 60]) and thin absorbers [57]. In general, close to ideal photocur-
rent collection is achievable whereas open circuit voltage and fill factor offer room for future
improvement. Figure 6.12 shows the external quantum efficiency of a Cu(In,Ga)S2 (E g ≈
1.53 eV) solar cell [61] which is close to unity at the maximum with a single layer antire-
flective coating. The curve is limited by the ZnO and absorber bandgaps at low and high
wavelengths, respectively. Holes generated in the CdS buffer layer are not collected which
leads to a drop in quantum efficiency for photon energies higher than the bandgap of CdS
(the interfaces on both sides are not inverted with respect to the n type CdS which leads to
a high recombination probability for holes, nevertheless, partial collection has been reported
   The bucking current is responsible for losses in open circuit voltage. It also leads to losses in
fill factor because the diode ideality factor A is higher than unity. In high efficiency CIGSe cells
the bucking current is due to bulk recombination in the space charge region of the absorber. The
diode ideality factor (between one and two) and its temperature dependence are in agreement
with analytical models describing recombination over an exponentially decaying density of


                       External quantum efficiency


                                                     0.4                        Cu(In,Ga)S2


                                                           400   500    600      700     800
                                                                   Wavelength (nm)

Figure 6.12 External quantum efficiency of a MgF2 /ZnO/CdS/Cu(In,Ga)S2 solar cell. The blue response
is limited by the bandgaps of ZnO and CdS, respectively. The red response is limited by the bandgap of
the absorber.

states [63, 64]. Such a defect distribution could be due to band tails and is also observed in
optical and admittance spectroscopy [65, 66]. Typical Urbach energies are in the range of 50
to 100 meV.
    In addition to the broad defect distribution there also seems to be a narrow defect dis-
tribution at 250–300 meV above the valence band [67]. It has been found by several meth-
ods and in samples prepared by various methods, including single crystals [25]. This de-
fect maintains its energetic position relative to the valence band within the whole CIGSe
system (for varied Ga/(Ga+In) ratios). However, its concentration seems to correlate with
the gallium content, with a minimum concentration found at Ga/(Ga+In) ≈ 0.3. Losses in
open circuit voltage were found to correlate with the concentration of this defect which sug-
gests its significance as a recombination path [68]. Its concentration is found to increase
upon radiation with energetic particles and to decrease with subsequent annealing at mod-
erate temperatures [69]. Other authors have identified a deeper defect at 0.8 eV above the
valence band edge [70]. This defect is approaching mid-gap position at high gallium con-
tents and can thus be expected to be an effective recombination center in wide gap CIGSe
    The saturation current in state-of-the-art low gap cells is mostly thermally activated and
the activation energy corresponds to the bandgap of the absorber. The diode ideality factor is
only mildly temperature dependent. Measurement of the open circuit voltage is as a function
of temperature and extrapolation yield: Voc (T = 0 K) = E g /q. Cells with inferior efficiency
show a stronger influence of tunneling. Depending on sample and temperature, higher and
more temperature dependent diode ideality factors are observed. If the tunneling influence
is not too severe, the bucking current mechanism can be described by tunneling assisted
recombination via trap states in the space charge region of the absorber [71]. In this case, the
activation energy of the saturation current still corresponds to the bandgap after taking into
account the temperature dependent diode ideality. Due to a common mistake in the evaluation
of measurement results [72] this model may have also been applied to devices where the
influence of tunneling is in fact much stronger.
                                               CHALCOPYRITE BASED SOLAR CELLS                       253

    Cells based on wide gap Cu(In,Ga)S2 absorbers exhibit open circuit voltages which are
lower than expected, considering the bandgap and the good bulk properties deduced from
photocurrent collection. Transport analysis [73] reveals that the dominant bucking current
mechanism changes with illumination. The ideal border cases are recombination over a re-
duced barrier at the interface without major assistance by tunneling (under illumination)
and recombination in the space charge region with significant tunnelling assistance (dark).
While this change of recombination mechanism with illumination is a unique feature of
Cu(Ga,In)S2 based cells [74], reduced thermal barrier and tunneling currents appear to be
more general problems and are also observed in wide gap C(I)GSe based solar cells [75,
76]. The achieved development status is illustrated by the exemplary efficiencies listed in
Table 6.1.
    The findings for the high efficiency cells described so far indicate that a further improvement
of efficiency beyond 20 % ought to be feasible by minimizing the bulk defect density. However,
it has also been argued that empirical device optimization has already led to an optimum
with respect to band edge fluctuations [77, 78]. Such fluctuations (which are also observed
in photoluminescence) may assist in charge separation, i.e. photocurrent collection, but are
detrimental in terms of bucking current. On the other hand, even if the efficiency of low gap
single junction cells is approaching practical limits, wide gap and tandem cells offer a large
potential for ultra high efficiency (Section 6.6.7).

Table 6.1 Efficiency of selected small area, laboratory style chalcopyrite based solar cells (partly with
antireflective coating)

Lab [Ref]                 Substrate    Absorber    Buffer       Efficiency    Remarks

NREL [79]                              CISe                       15 %       E g = 1.04 eV
NREL [80]                              CIGSe                      19.5 %     E g = 1.15 eV
NREL [79]                              CIGSe                      10.2 %     E g = 1.64 eV
NREL [79]                 glass        CGSe        CdS             9.5 %     E g = 1.68 eV
HMI [17]                               CIS                        11.4 %     E g = 1.5 eV
HMI [61]                               CIGS                       12.3 %     E g = 1.53 eV
HMI [61]                               CIGS                       10.1 %     E g = 1.65 eV
NREL [81]                 steel                                   17.5 %     flexible
HMI [82]                  titanium     CIGSe       CdS            16.2 %     flexible
ETHZ [83]                 polyimide                               14.1 %     flexible
NREL/Aoyama               glass        CIGSe       Zn(S,O,OH)     18.6 %     Cd freea buffer by CBD
  Gakuin Univ. [84]
Shell Solar/HMI [85]                   CIGSSe      In2 S3         14.7 %     Cd free buffer by ILGAR
Shell Solar/HMI [86]                   CIGSSe      (Zn,Mg)O       12.5 %     Cd free buffer by
                                                                               sputtering (dry process)
ZSW/CNRS [87]                          CIGSe       In2 S3         16.4 %     Cd-free buffer by
                                                                               ALCVD (dry process)
Stuttgart Univ. [88]                   CIGSe       Inx S          14.8 %     Cd-free buffer by
                                                                               evaporation (dry

      see [89] for a more complete list of Cd free devices.

Figure 6.13 Partial view of the pilot production line for CIS modules (substrate size 120 × 60 cm2 )
at Sulfurcell in Berlin, Germany. Sulfurcell uses the process sequence shown in Figure 6.3. Reproduced
with the permission of Dr. Nikolaus Meyer, Hahn-Meitner-Institut Berlin GmbH.

Efficient chalcopyrite based photovoltaic modules were demonstrated years ago [90]. Despite
this early proof-of-concept, the status has been limited to pilot lines (Figures 6.13 and 6.14),
or announcements of such, for a long time. Consequently, chalcopyrites do still not play a
significant role in the marketplace. Nevertheless, industrial laboratories have always had a
major impact on development and an increasing number of companies world wide is involved
in the development of commercial products for the power market as well as niche applications.
The first commercial products, announced by Siemens Solar Industries in 1998, were small

Figure 6.14 Large area in-line sputter coater (Von Ardenne Anlagentechnik, Dresden, Germany) for
the deposition of molybdenum and zinc oxide at the Wuerth Solar pilot line in Marbach, Germany.
Reproduced with the permission of Von Ardenne Anlagebtechnick GmbH.
                                          CHALCOPYRITE BASED SOLAR CELLS                       255

modules with 5 and 10 W rated output power manufactured on a pilot line in Camarillo, CA.
The biggest producers today are Shell Solar and Wuerth Solar. Shell Solar is now operating the
production line in Camarillo which has been upgraded with respect to production capacity and
product size. The largest Shell Solar modules are rated at 40 W, the Wuerth Solar modules go
up to 80 W due to their larger panel size (60 × 120 cm2 ). Both, Shell Solar and Wuerth Solar
have also published yield data and efficiency distribution of their (pilot) production which
clearly demonstrate the feasibility of mass production of chalcopyrite based modules. Wuerth
Solar has announced a new factory with a rated production volume of 15 MWp/a from 2007
onwards [91].
   Nevertheless, the transition from laboratory to large scale manufacturing has, in general,
been more difficult than expected. According to [92] the main obstacles have been:

r Commercial scale equipment.
r Quality control and in situ monitoring.
r Uniformity.
r Low open circuit voltage.
r Throughput.
r Stability.

    For some process steps the machines for large area deposition had to be custom designed
based on experience with lab scale equipment which is an expensive and error prone pro-
cess. With the increasing number of installations the equipment manufacturers are gaining
experience in manufacturing these systems. The rapid commercialization of other thin film
technologies provides synergy effects. Glass companies experienced in very large area thin
film coating of glass (low emission coating) are using this experience to offer glass substrates
already coated with diffusion blocker and molybdenum. Transparent conductive oxide sput-
tering machines (Figure 6.14) can be derived from those sold for preparation of TCO coatings
for flat panel liquid crystal and organic displays. Equipment manufacturers are now actively
participating in the establishment of pilot production lines and it is to be expected that turn-key
production facilities with process and product specification will be available in the not too
distant future.
    Criteria for go/no go classification after each individual process step and methods for quality
control by in situ monitoring had been neglected for a long time in basic research. However,
significant progress has been made in recent years as more research laboratories have also
become involved in scaling up activities. Examples include Raman spectroscopy [93] and
photoluminescence decay [94] to assess the absorber quality, and X-ray fluorescence [95]
for in system stoichiometry measurements. Laser light scattering and substrate temperature
monitoring have already been described in Section 6.3.1.
    Uniformity or low open circuit voltage appear to be no longer a general problem as indicated
by the module efficiencies that have been achieved in practice (Table 6.2).
    Throughput in sequential processes has been improved by the development of the rapid
thermal annealing processes. For evaporation technology a feasible approach appears to be to
scale up the sources for increased width of deposition. Further throughput improvements are
desirable for large scale production. Minimizing the required thickness of individual layers

Table 6.2 Efficiency of selected minimodules and full size modules

Company or Institute           Size               Efficiency      Remarks

Hahn-Meitner-Institut [17]     5 × 5 cm2            9.7 %        Se free CIS
Uppsala University [4]         5 × 5 cm2           16.6 %        With grid
Showa Shell [5]                30 × 30 cm2         14.2 %        Cd free
Shell Solar [6]                60 × 90 cm2         13.1 %
Global Solar Energy [108]      7085 cm2            10.1 %        Not monolithically interconnected,
                                                                   cells on metal foil
Wuerth Solar [7]               60 × 120 cm2        12.2 %

[96], alternative methods for patterning, eliminating the buffer layer from the module structure
[86], and high rate reactive TCO sputtering carry potential for such improvement.
    Outdoor testing of modules has generally demonstrated excellent stability [98, 99]. Owing
to the increasing production volume there is a growing number of installations (Figure 6.15)
where the actual performance [100] and long term stability can be assessed (Figure 6.16:).
Accelerated lifetime testing, especially the damp heat testing procedure which forms a part
of the EN/IEC 61646 certification, has, however, been cumbersome [101, 102]. Partly, this
is due to transient effects which occur during stress tests. These can lead to an apparent
degradation, however, the efficiency recovers after several days of light soaking. The exact

Figure 6.15 Solar Tower, Training and Technology Center Handwerkskammer Heilbronn, Germany.
Wuerth Solar frameless CIS Fa¸ ade modules with a total nominal power of 8 kWp (STC) were installed
in April 2001. The installation comprises 120 modules with the dimensions 60 cm by 120 cm and 40
modules with the dimensions 40 cm by 60 cm. Reproduced with permission of Wurth Solar GmbH &
                                           CHALCOPYRITE BASED SOLAR CELLS                        257

Figure 6.16 Results of long term outdoors measurements of CIGSSe based module prototypes. Modules
installed at the National Renewable Energy Lab (NREL) outdoor test facility (OTF) in Golden, Colorado
[97]. Reproduced with permission of Shell Solar Industries.

causes for degradation are still under investigation but empirical optimization has already
achieved modules which have been independently certified [6, 99].
   In conclusion, technical problems have been overcome to a large extent. On the other
hand, they still contribute to the financial risk. As long as specialized deposition equipment
is produced only in small numbers and turn-key facilities are not available, the technology
development has to be taken step by step. This implies a considerable lead time before the
production volumes are high enough to achieve competitive production costs in relation to
the conventional silicon modules produced in large scale facilities. Innovative niche market
products which exploit inherent features of chalcopyrite thin film devices not readily available
with silicon can be sold at higher prices. While they can offer a faster return on investment
for smaller (start up) companies, the true medium and long term benefit, as acknowledged
by independent studies, lies in the cost reduction potential exploitable only through mass
production for the power market.

6.5.1 Cost estimations
An early cost study [103] predicted that chalcoypyrite based modules could be manufactured
at 0.6 € /Wp in a plant with 60 MWp annual capacity whereas even the most cost effective
multicrystalline silicon based technology would require a 500 MWp/a factory to achieve sim-
ilar costs [104, 105]. Another study [106] comparing the direct module manufacturing costs
of single and multicrystalline silicon as well as amorphous silicon, CdTe, and chalcopyrites
estimated the cost to be 2.25 $/Wp for a 10 MWp/a chalcopyrite production line. This was
lower than the cost for all other technologies, except for multicrystalline silicon estimated at
2.10 $/Wp. The estimation was based on the assumption of 9 % average module efficiency and
65 % production yield for the chalcopyrite modules, numbers that are clearly more favourable in

today’s pilot production. It was estimated that a 100 MWp/a production capacity at improved
average efficiency and yield could result in costs of 1 $/Wp, 15 % lower than that of mul-
ticrystalline silicon. In general, the study was based on a first generation baseline process and
additional cost benefits are expected from current technology updates. A more recent estima-
tion for CuInS2 modules [107] already claims manufacturing costs of 1.5 €/Wp for a small
production line (5 MWp/a).

6.5.2 Module performance
It has been pointed out already that chalcopyrite cells offer a very high efficiency potential.
Monolithic integration and large area nonuniformity cause only small losses, therefore high
efficiencies could also be demonstrated for test structures (minimodules) and even full size
modules (Table 6.2).
    Considering mass production, the distribution of efficiencies (or power output for a given
module size) is much more relevant than the record efficiency of a single module. Wuerth Solar
reports [99] an output of 79.9 ± 2.2 Wp for a batch of 306 modules (60 × 120 cm2 ) indicating
that a very narrow distribution curve is feasible. In 2004 their production yield was more
than 80 % with an average module efficiency slightly higher than 11 % [109]. In 2002 Shell
Solar reported [110] an average efficiency of 10.9 % for nearly 16 000 laminates produced in
Camarillo (Figure 6.17).
    Considering the application the rated power output is an important point. However, the
module has to perform well not only at the standard testing conditions applied for measur-
ing the rated power but also at conditions typically encountered in operation, such as lower
illumination intensity, varying spectral distribution, partial shading and elevated temperatures
(performance ratio). Chalcopyrite modules have shown high power output under these con-
ditions. The typical interconnection scheme with a large number of narrow cells reduces the

Figure 6.17 Distribution of 1 × 4 laminates produced in by Shell Solar in 2002. Data includes some
15 785 laminates. Gaussian fit has an average of 10.9 ± 0.6 %. Shown for comparison are the lower
specification limits for the product family [110]. Reproduced with permission of Shell Solar Industries.
                                           CHALCOPYRITE BASED SOLAR CELLS                       259

impact of partial shading. Shunts have been reduced resulting in better performance at low
illumination intensity. Higher absorber bandgaps limit the losses with increasing module tem-
perature [100]. Consequently, it has been observed that chalcopyrite modules can outperform
silicon modules in terms of the annual energy output on a kWh/kWp basis [99, 109, 111].

6.5.3 Sustainability

Sustainability has many aspects and while not all relevant issues are completely clarified at
this point the available data are quite promising (see below). Due to the high efficiency and
performance ratio, the long lifetime, the low material consumption, and low thermal budget
a chalcopyrite module will have a favourable energy balance. There are a number of research
efforts [112] addressing specifically the sustainability of mass production. Like any large scale
deployment, mass production of chalcopyrite PV modules will result in production related
waste materials, energy consumption, and raw material depletion. Recycling of waste and
modules at the end of their lifetime is mandatory for a sustainable technology. In addition, the
producer may at some time be legally required to take back modules as is the case already for
electronic products in Europe [113]. Availability of raw materials and recycling

Thin film technologies make very efficient use of raw materials. While 0.5–1 kg/m2 of semi-
conductor grade silicon are required for a conventional module, the material consumption per
square meter for the active films of a CISe module is given as: 7–20 g molybdenum, 1.5–4 g
copper, 3–9 g indium, 7–20 g selenium, and 1–3 g zinc (depending on the exact module struc-
ture and yield [114]). This implies that the total material input is comparable to the material
used for just the grid metallization of silicon modules.
   Nevertheless, it has been argued that indium is a bottleneck concerning the abundance of
raw materials. In 2003 it was used mainly for coatings (65 %), solders and alloys (15 %), and
electrical components (10 %) [115]. Indium based coatings are used in the production of flat
panel displays where indium tin oxide (ITO) is used as a transparent contact. The annual world
production of indium, mainly from zinc ores, is in the order of 300 tons, which translates into
chalcopyrite modules with approximately 15 GWp/a in total. On the other hand, indium is
about three times more abundant in the Earth’s crust than silver, the latter having an annual
production of 20 000 tons and a reserve base of 570 000 tons [115]. These numbers imply that
the availability of indium is not likely to be an ultimate limiting factor. Moreover, it has already
been shown that even thinner active films in a chalcopyrite module are feasible [116]. The flat
panel industry may replace ITO by the cheaper ZnO in future and indium recycling will, in
addition, contribute to higher availability and lower market prices. Indium free absorbers are
under development (see Section 6.6.3). Within the last three decades the prices for indium
have been varying over a wide range from below 100 $/Kg to more than 500 $/Kg. However,
even the latter price would imply that indium is responsible for only approximately 2 % of the
module manufacturing costs.
   Waste from the dry processes consists of material that is deposited onto chamber walls,
shutters, substrate carriers etc. Recycling should be possible, but the amounts are probably too
low to make it attractive unless mass production has started. Sputter targets would presumably

be returned to the manufacturer to reclaim the remaining raw materials. Research has put more
emphasis on recycling the waste generated by wet chemical buffer layer deposition [117, 118].
This could be partly done directly at the production site by removing reaction products from
the solution then readjusting the concentrations and feeding the solution back to the process.
    A complete process sequence for disassembling and recycling of off spec modules and
semiproducts has been tested successfully and could also be used for end-of-life modules
[119]. The module is heated to 250 ◦ C which softens the EVA encapsulation and allows a
mechanical removal of the cover glass. Window and buffer layers are etched away by a mild
acidic solution. The chalcopyrite film is scraped from the back contact and finally the latter
is dissolved in nitric acid. The test confirms the feasibility of disassembling the module layer
by layer which results in low cross contamination. It has been suggested that these relatively
clean materials, such as the chalcopyrite powder, could be used directly in adapted preparation
processes. It has already been shown that the substrate glass can simply be reused for new
modules. Energy payback time

A comparative study on energy payback times has been assisted by the only company that
produces silicon as well as chalcopyrite modules [120]. It was concluded that the EPBT for
the chalcopyrite module is 1.8 years as compared to 3.3 years for the silicon module. It should
be noted that the aluminium frame is responsible for a large fraction of the materials energy
content of the chalcopyrite module. The study was also based on first generation chalcopyrite
technology and significantly lower EPBT should be feasible.

The chalcopyrite module is still under active international development. Progress in fundamen-
tal understanding and preparation technology will result in significant improvements in market
potential, module performance, sustainability, and minimized ecological impact of large scale
production of next generation modules. We will highlight areas of ongoing development in the
following paragraphs.

6.6.1 Lightweight and flexible substrates
Transferring the technology of chalcopyrite based solar cells from rigid glass substrates to
flexible, low-mass substrates [26] opens new market segments. Furthermore, flexible substrates
are a requirement for roll-to-roll processing which could lower production costs [108]. Such
substrates can be plastic or metal foils (Figure 6.18). When combined with foil substrates,
the low mass, excellent radiation hardness, and relatively high (compared to other thin film
technologies) efficiency make chalcopyrite based technology the leading candidate for thin
film cells for space applications [121].
   Available plastic foils do not tolerate the high substrate temperatures ordinarily used for
chalcopyrite preparation. Common problems at high temperatures are shrinkage, warping,
outgassing, and loss of flexibility. Lowering the substrate temperature is possible, but the
                                            CHALCOPYRITE BASED SOLAR CELLS                           261

Figure 6.18   Prototype of a flexible CIGSe solar cell for space applications on a thin titanium substrate

efficiencies that can be achieved are, in consequence, somewhat lower. The best results, still
below 15 % efficiency, are achieved on polyimide foils [83]. On the other hand, the conductivity
of metal foils raises new challenges with respect to pin hole free processing and monolithic
integration. Some metals also tend to diffuse into the absorber and cause deterioration of
its properties. Isolating and/or diffusion blocking layers are under development to circumvent
these problems. In another approach monolithic integration is abandoned altogether. Individual
medium sized cells are fabricated and interconnected using a shingling scheme.
    Sodium doping is mandatory with any foil substrate and can be carried our through a
precursor film, coevaporation or even diffusion after film preparation.

6.6.2 Cadmium free cells
The CdS buffer layer is very thin and contamination of the environment from modules is very
unlikely even in extreme conditions. Nevertheless, cadmium is a hazardous material and its
elimination from the module may increase the general acceptance of the product and reduce
production costs (by avoiding costly safety measures). According to European Union directives
[122] several heavy metals (including cadmium) must be not be contained in new electrical and
electronic equipment after July 2006. While photovoltaic modules do not currently fall under
these regulations [113] they illustrate the general effort to reduce the amount of cadmium in
circulation. Owing to the intensive work on cadmium free buffer layers there is a variety of
possible alternatives. We will introduce a small selection in this chapter. More information can
be found in a recent review article [89]. Wet chemical processes

Chemical bath deposition is not restricted to the deposition of CdS. The most common re-
placements are indium or zinc based sulfides, oxides, hydroxides, or mixtures thereof. The

achievable efficiency is very close to that of the standard device [84]. However, the deposition
parameters appear to be more critical and need to be adapted to the specific absorber. Accord-
ingly, it can be challenging to achieve a high reproducibility of cell performance. Cadmium
free cells are often found to exhibit metastability effects, e.g., significant improvement with
light soaking [123]. It is not yet fully clear how far the long term stability of the cell is also
affected by the modified buffer layer. Only one of the current pilot production efforts is known
to use a zinc based buffer layer (Table 6.2).
   Novel chemical deposition methods such as ILGAR (ion layer gas reaction) have been
developed and have been applied successfully mostly to those Cd free alternative buffer layers
[124]. There are indications that ZnO deposited by ILGAR can combine the functions of the
buffer and the (normally sputtered) undoped part of the window layer [125]. Dry processes

Wet chemical deposition is often believed to be significantly cheaper than other deposition
methods. However, in a typical processing sequence, buffer layer deposition is the only wet
chemical process, which implies that the whole infrastructure has to be implemented for just
this single process step. In addition, it prevents true in-line processing. While this is currently
not a big issue because the deposition systems are typically not connected and scribing is also
carried out in atmosphere, it may be a drawback in view of future large volume production.
Dry processes are therefore under investigation. The most successful approaches have been
MOCVD [126] and ALCVD (atomic layer chemical vapor deposition [127]). The former has
so far produced only small samples, albeit with good efficiency, whereas the latter has been
shown to work at least on medium sized substrates (30 × 30 cm2 ). Atomic layer chemical
vapor deposition is, however, an inherently slow process and therefore only cost-effective if
carried out in systems that can accommodate a large number of substrates simultaneously, in
contradiction to the requirements for an in-line capable process.
   Buffer layers can also be evaporated which appears particularly attractive if the absorber is
also prepared by evaporation [128]. Efficiencies of almost 15 % can be achieved provided the
buffer layer is evaporated insitu, i.e., without breaking vacuum after the absorber deposition
   Evaporation, MOCVD, and ALCVD, in analogy to CBD, are ‘soft’ deposition methods and
there have been assumptions that this is a major advantage. On the other hand, sputtering of a
Cd free buffer layer, if feasible, is clearly a technically very attractive solution. Sputtering is,
however, not a soft method and it has been speculated that one of the main functions of the buffer
layer consists of protecting the absorber surface from sputter damage during window layer
deposition. This assumption is supported by the typically poor device performance observed
when simply omitting the buffer layer and sputtering the TCO directly onto the absorber. It
was shown, however, that a chemical treatment (partial electrolyte) is, in principle, sufficient
to stabilize the absorber surface for window layer sputtering [129]. Only a very thin film is
deposited under the conditions of this chemical treatment, which would not be sufficient to
prevent sputter damage. More recently it has been shown that an efficiency of more than 12 %
can be achieved by a sputtered Inx S buffer layer without chemical treatment [130].
   A different novel approach to heterojunction formation originates from the question whether
the window layer can be modified in such a way that a buffer layer is no longer necessary. Based
on the assumption that the band line up at the chalcopyrite/TCO direct junction is causing the
                                         CHALCOPYRITE BASED SOLAR CELLS                      263

poor performance, experiments have been carried out using (Zn,Mg)O rather than pure ZnO in
the first part of the TCO double layer stack for improved band alignment. Kelvin probe force
microscopy measurements of the work function of solar cell cross sections support this model
[34]. In terms of cell efficiency (12.5 %), the proof-of-concept could be achieved [131], but
further development is necessary to demonstrate long term stability and reproducibility in pilot
production. It has been shown that the alloy can be deposited by sputtering from a single mixed
target and that its conductivity is comparable to the standard undoped ZnO. Hence, sputtering
of (Zn,Mg)O is a simple drop-in replacement and can be integrated easily into existing pilot
lines to prepare standard (with CBD buffer) modules and to investigate modules without buffer
layers in the same setup.

6.6.3 Indium free absorbers
It is sometimes argued that the lack of abundancy of indium may be limiting the long term per-
spectives of chalcopyrite based photovoltaics (see Section This has triggered research
on In-free absorbers. Compounds such as Cu2 ZnSn(S,Se)4 crystallize in the kesterite structure
which can be derived from the chalcopyrite structure by replacing half of the In atoms by Zn
and the other half by Sn atoms. It has been shown that the synthesis of crystals and thin films
can be carried out using methods well known from chalcopyrite preparation [132, 133]. Also,
heterojunction solar cells can be prepared using the established contact layers (molybdenum
back contact, CdS buffer etc.). On the other hand, secondary phases are more problematic in
this quarternary system as compared to ternary chalcopyrite. Cell efficiencies achieved so far
are in the range of 5 % [134] and indicate that research efforts would have to be stepped up
significantly to make kesterites a feasible alternative to chalcopyrite absorbers.

6.6.4 Novel back contacts
Molybdenum appears to be an almost ideal contact material at least for current typical cell
structures and preparation methods. Corrosion of molybdenum [135] is one of the few known
issues and has been reported to contribute to module degradation in accelerated lifetime testing.
It is conceivable that the stability could be improved by using molybdenum based alloys instead
of pure molybdenum. Another disadvantage of molybdenum is its poor optical reflection which
may become relevant in view of efforts to reduce the absorber thickness. The choice of other
metals is limited due to their instability under typical absorber deposition conditions. Tungsten
seems to yield a good ohmic contact but its optical properties in this respect are poor. Tantalum
and niobium have slightly higher reflections and preliminary studies [136] indicate that they
may be feasible in terms of contact performance. Transparent conductive oxide coated metal
contacts may be an alternative solution to achieve stability, and good electrical as well as
optical performance for thin cells. Ohmic TCO/chalcopyrite contacts are also required for
novel structures such as bifacial and (semi) transparent cells (see below).

6.6.5 Bifacial cells and superstrate cells
In a variant of the cell structure the nontransparent rear metal contact is replaced by a TCO film
[137]. Such a bifacial cell (Figure 6.19) has advantages in certain applications. At the current

               Substrate                  Superstrate               Bi-facial

                                                                    Blocking Contact
                                                                    Ohmic Contact
                TCO    Chalcopyrite Glass         Metal

       Figure 6.19   Schematic cross sectional views of chalcopyrite solar cell configurations.

development state the efficiency for illumination through the back contact is significantly lower
because the carriers are generated outside the field zone in proximity to the poorly passivated
contact (poor blue response). Further optimization of absorber thickness, diffusion length, and
contact passivation appears to be feasible.
    In the conventional module structure the nontransparent back contact is deposited onto
the glass substrate. The module is illuminated through the encapsulation material which has
to be transparent (usually a second sheet of glass). It is, in principle, possible to reverse the
cell structure by starting with the deposition of the transparent contact (Superstrate configura-
tion, Figure 6.19). In this case the TCO/(buffer)/chalcopyrite interface needs to be a blocking
junction. The light enters the cell through the superstrate which has the advantage that the
module can be encapsulated with nontransparent material of lower mass and lower cost. The
efficiency in this structure is reduced due to junction degradation during the high temperature
absorber preparation and/or inferior quality of the absorber deposited at lower temperatures.
The proof-of-concept is limited to small area cells [138–140].
    Heterojunctions can also be prepared on TiO2 films [141]. This material can be grown so
that it is nanostructured (porous) and is better known from its application in dye sensitized
cells. Filling the pores of TiO2 with a chalcopyrite semiconductor has been proposed as a
realization of the ETA (extremely thin absorber) concept [142]. The recombination velocity
at the TiO2 /chalcopyrite interface must be kept low in view of the very large interface area in
this structure. It has been found that blocking or buffer layers are useful in achieving better
interface properties [143, 144].

6.6.6 Nonvacuum processing
Nonvacuum processing is being investigated because it may reduce the up front investment
required to implement a certain process step. This lowers the initial barrier and financial risk
but does not necessarily imply lower overall production costs. The running costs may be
actually be higher due to the mix of vacuum and nonvacuum equipment and corresponding
                                          CHALCOPYRITE BASED SOLAR CELLS                      265

infrastructure requirements, larger amounts of waste generated in wet chemical processing,
etc. There is no known vacuum-free alternative for the preparation of the molybdenum back
contact, however, metallized glass may be bought from the glass industry and does not have
to be prepared on site. Nonvacuum buffer (CBD) and window layer (MOCVD) deposition
is already established. Nonvacuum absorber processes are commonly based on sequential
processes similar to those described in Section above. They comprise a precursor
deposition at room temperature followed by (reactive) annealing at elevated temperature. This
second stage of sequential processing, i.e., selenization in mixtures of H2 Se and inert gas is an
already established nonvacuum process. Electrodeposited metal precursors were investigated
early in the history of chalcopyrite solar cells [145] and there appears to be a renewed interest
[146] in spite of the moderate lab scale efficiencies achieved. Electrodeposition of a metal/Se
precursor is the basis of a process developed with strong industrial participation and so far
yielding small area cells with 10 % efficiency [147]. Higher efficiencies could be achieved by
another group with selenium containing precursors, but only after manipulating the precursor
composition by means of conventional evaporation [148].
   Another interesting approach is the preparation of inks or slurries which can be deposited
by spraying, printing, dip coating, or doctor blading. In one approach [149] the metals are
dissolved in acid and hydroxide nanoparticles are precipitated from this solution. The particles
are dried, which results in a fine powder of mixed oxides. The powder is dispersed in an
aqueous solution to obtain the ink. In this and similar processes the film composition can be
controlled precisely and with relative ease by adjusting the relative amounts of material used
for particle preparation. The oxide particle precursor needs an additional reduction step prior
to selenization carried out in diluted hydrogen. The precursor film can be porous because the
volume increase during selenization leads to sufficient densification [150]. In principle it would
be possible to include selenium (and/or sulfur) in the precursor particles and to use nonreactive
sintering, but results from various approaches were not convincing.

6.6.7 Wide gap and tandem cells
Wide gap cells are interesting for two reasons: better single junction cells and top cells for
tandem configuration. A moderate increase of the bandgap to about 1.4 eV for a single junc-
tion cell places it at the theoretical maximum of the bandgap/efficiency relation for the solar
spectrum [151]. Because of the reduced current density, a higher resistance of the TCO can be
tolerated. At the same time the doping of the TCO can be at the upper limit because the absorber
cut off wavelength is below the onset of free carrier absorption in the TCO. In consequence the
TCO can be made thinner which leads to cost reductions. Compared to low gap absorbers, the
loss of performance at higher module temperature is much less severe which leads to improved
annual output in terms of kWh/kWp (performance ratio) especially in hot climate. Tandem
cells could eventually lead to an efficiency exceeding even the theoretical limit for silicon
cells. Low gap CIG(S)Se cells (E g ≈ 1.1 eV), such as NREL’s 19.5 % efficient cell, are ideally
suited as the bottom cell. The top cell absorber should have a bandgap of roughly 1.7 eV [152]
and the front as well as the back contacts need to be highly transparent.
   Ideally, the top cell would be deposited directly onto the bottom cell (monolithic tandem,
Figure 6.20). This, however, requires fundamental changes to the chalcopyrite technology. The
top cell, after laying down the window layer, does not tolerate temperatures above 200 ◦ C. This
maximum allowable temperature would have to be raised significantly and low temperature

                                                   Mechanically Stacked


                               Low gap Wide gap

                       TCO         Chalcopyrite     Glass     Metal     EVA

                                                  Blocking Contact
                                                  Ohmic Contact

              Figure 6.20    Schematic cross sectional views of tandem configurations.

processing would have to be implemented for the deposition of the top cell. It is therefore more
likely that the first tandem modules will be mechanically stacked after individual preparation
of top and bottom cells, respectively. If the top module were prepared in analogy to the bifacial
cell described above this would imply a laminate with three sheets of glass. If the top module
had a superstrate structure, top and bottom could be laminated together face-to-face and, like
in the conventional single junction module, only two sheets of glass would be necessary for
an encapsulated module (Figure 6.20, right hand side).
    In practice the realization of highly efficient wide gap cells has not yet been achieved. As
pointed out previously, this may have partly historical reasons. On the other hand, significant
differences in parameters determining the cell efficiency have been theoretically predicted
[153] and are observed in the experiment [154]. It is not fully clear which of these are the
most crucial and how they could be resolved. This is expected to change as the development of
widegap cells and tandems is one of the focus points of international basic research. Preliminary
prototypes of monolithic [155] as well as mechanically stacked [156] tandems are described
in the literature. Due to the problems in realizing a highly efficient wide gap chalcopyrite solar
cells it has been suggested to use a different material such as amorphous silicon [157] or CdTe
[158] for the top cell.

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