Solar Energy Materials & Solar Cells 95 (2011) 1452–1462

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Solar Energy Materials & Solar Cells

0927-02

doi:10.1

n Corr

E-m

journal homepage: www.elsevier.com/locate/solmat

Analysis of Cu(In,Ga)(S,Se)2 thin-film solar cells by means ofelectron microscopy

D. Abou-Ras a,n, J. Dietrich b, J. Kavalakkatt a, M. Nichterwitz a, S.S. Schmidt a, C.T. Koch c, R. Caballero a,J. Klaer a, T. Rissom a

a Helmholtz-Zentrum Berin fur Materialien und Energie, Hahn-Meitner-Platz 1, 14109 Berlin, Germanyb Department of Semiconductor Devices, TU Berlin, Einsteinufer 19, Sekr. E2, 10587 Berlin, Germanyc Max-Planck Institute for Metals Research, Heisenbergstr. 3, 70569 Stuttgart, Germany

a r t i c l e i n f o

Available online 16 December 2010

Keywords:

Cu(In,Ga)(S,Se)2 solar cells

Electron microscopy

Electron backscatter diffraction

Electron-beam-induced current

Electron energy-loss spectroscopy

Electron holography

48/$ - see front matter & 2010 Elsevier B.V. A

016/j.solmat.2010.11.008

esponding author.

ail address: daniel.abou-ras@helmholtz-berlin

a b s t r a c t

The present work gives an overview of how electron microscopy and its related techniques are used to

analyze individual layers and their interfaces in Cu(In,Ga)(S,Se)2 thin-film solar cells. Imaging of samples

can be performed at scales of down to the (sub)angstroms range. At similar spatial resolutions,

information on composition can be gathered by means of energy-dispersive X-ray spectroscopy (EDX)

and on spatial distributions of electrostatic Coulomb potentials in the specimen by applying electron

holography. Microstructural and compositional properties as well as charge-carrier collection and

radiative recombination behavior of the individual layers are accessible by use of electron backscatter

diffraction, EDX, electron-beam-induced current (EBIC) and cathodoluminescence measurements,

available in scanning electron microscopy. The present contribution gives an overview of the various

scanning and transmission electron microscopy techniques applied on Cu(In,Ga)(S,Se)2 thin-film solar

cells, examples from case studies, and also demonstrates how these techniques may be combined in order

to improve the analysis. Particularly, EBIC results show a reduced charge-carrier collection at Cu(In,Ga)Se2

grain boundaries, while no indication was found for a charge accumulation at the grain boundaries by

electron holography.

& 2010 Elsevier B.V. All rights reserved.

1. Introduction

Electron microscopy analyses have been performed onCu(In,Ga)(S,Se)2 thin-film solar cells ever since the beginning oftheir research and development more than 30 years ago. Apart fromproviding images down to the angstroms range of the thin-filmstacks with total thicknesses of few micrometers, electron micro-scopy supplies a versatile collection of various analysis methods forstudying microstructures and compositions as well as of electricaland optoelectronic properties.

The present work intends to give an overview of some of thesemethods, describing how they are applied to provide insight intothe physics of Cu(In,Ga)(S,Se)2 thin-film solar cells. Particularly,the benefit of the combination of various techniques onidentical positions is highlighted when studying structuraldefects and compositional gradients or also grain boundaries inCu(In,Ga)(S,Se)2 thin-film solar cells.

ll rights reserved.

.de (D. Abou-Ras).

2. Experimental details

Co-evaporated Cu(In,Ga)Se2 and CuInS2 produced by a rapidthermal process (RTP) were deposited on Mo-coated soda-limeglass substrates. For the Cu(In,Ga)Se2 layers, a multi-stageprocess [1] was applied, in which first In–Ga–Se is evaporated,then Cu–Se until the layer exhibits a concentration ratioof [Cu]/([In]+[Ga])41, and finally again In–Ga–Se until[Cu]/([In]+[Ga])o1. In the RTP [2], Cu–In metal precursors weresputtered on Mo/glass and then transformed into CuInS2 in a sulfuratmosphere. Since the Cu concentration exceeded the one ofIn ([Cu]/[In]41) in the precursors, Cu–S formed on top of CuInS2,which was removed later by use of diluted KCN, prior to furtherprocessing.

The solar cells were completed by chemical-bath-deposited CdSbuffers, a sputtered i-ZnO/ZnO:Al bilayer as window layers and aNi–Al metal grids for enhancing the current collection. Table 1 givesan overview of the samples studied for the present work, includingtheir compositions and photovoltaic parameters.

Cross-section specimens for scanning electron microscopy(SEM) were prepared by gluing two stripes of solar cells face-to-face together using epoxy glue, cutting stripes from these stacks,then by polishing the cross-sections of these stripes mechanically

Table 1Compositional ratios and photovoltaic parameters (open-circuit voltage Voc, short-circuit current density jsc, fill factor FF, and solar conversion efficiency Z) of Cu(In,Ga)Se2

(CIGSe), CuGaSe2 (CGSe), and CuInS2 (CIS) thin-film solar cells studied for the present work.

Sample [Cu]/([Ga]+[In]) [Ga]/([Ga]+[In]) Voc (mV) jsc (mA/cm2) FF (%) Z (%)

CIGSe-1 0.86 0.33 633 34.3 70 15.2

CIGSe-2 0.92 0.33 700 31.0 76 16.4

CIGSe-3 0.83 0.28 674 31.1 71 14.8

CIGSe-4 0.86 0.37 704 32.6 72 16.6

CGSe-1 0.80 1 702 13.1 55 5.1

CIS-1 1 – 692 21.4 68 10.1

CIS-2 1 – 694 20.8 66 10.1

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–1462 1453

and using Ar-ion beams. Very thin (4–5 nm nominally) graphitelayers were deposited on top of these cross-sections in order toreduce drift during the measurements. Backside Cu(In,Ga)Se2

layers were exposed by lifting off Cu(In,Ga)Se2/CdS/i-ZnO/ZnO:Alstacks from the Mo/glass substrates, after the stacks had been gluedto a SEM holder by use of Ag epoxy glue. Very thin graphite layerswere deposited on the ZnO:Al layers prior to gluing and also on theexposed Cu(In,Ga)Se2 layers in order to inhibit Ag diffusion into thesolar-cell stack and to reduce drift during the measurements.

Cross-section specimens for transmission electron microscopy(TEM) were prepared by gluing two stripes of solar cells face-to-face using epoxy glue, cutting stripes from these stacks, then bypolishing the cross-sections of these stripes mechanically andgluing them to Mo rings. The polished cross-sections were Ar-ion-milled until they were transparent for the electron beam.

Electron-backscatter diffraction (EBSD), energy-dispersive X-ray spectrometry (EDX), electron-beam-induced current (EBIC) andcathodoluminescence (CL) measurements were performed using aLEO 1530 GEMINI scanning electron microscope equipped with anOxford Instruments HKL Nordlys II EBSD camera (acquisition andevaluation software FastAcquistion/Channel5), a Thermo Noran X-ray silicon-drift detector (acquisition and evaluation softwareNoran System Seven), a FEMTO DLPCA-200 impedance amplifier,and an EMSystems CL system (acquisition software Proscan SpectraInterpreter) with a Hamamatsu photomultiplier tube R3896.

TEM images and EDX maps were acquired using a Zeiss Libra200FE transmission electron microscope. Details on the recon-struction of the amplitude and phase of the exit-plane wavefunction from focus series acquired by using the same microscope,also called inline electron holography, can be found in Refs. [3,4].

3. Results and discussions

3.1. Analyses of microstructures, compositions and optoelectronic

properties

3.1.1. Scanning electron microscopy

3.1.1.1. Electron backscatter diffraction. A straightforward methodin order to obtain information of microstructural details of aCu(In,Ga)(S,Se)2 thin film, such as its average grain size, is tofracture the solar cell (particularly for those on glass substrates)and image the cross-section. The average grain size, e.g., may beestimated from such images under the assumption that the thinfilm fractures along the grain boundaries, which needs not to bethe case.

A more precise method for measuring average grain sizes iselectron-backscatter diffraction [5,6] on polished cross-sectionspecimens or on chemically etched surfaces of thin films. Suchsurfaces of Cu(In,Ga)(S,Se)2 layers may also be exposed by lifting offCu(In,Ga)(S,Se)2/CdS/ZnO stacks from the Mo/glass substrates. Thecareful surface preparation of specimens for EBSD is importantsince this technique is surface sensitive with information depths in

Cu(In,Ga)(S,Se)2 of only about 20–30 nm. Fractured cross-sectionsamples allow for pointwise EBSD measurements, but contiguousEBSD maps are impeded by shadowing of emitted backscatteredelectrons by the surface roughnesses. EBSD is based on thediffraction of backscattered electrons at various families of latticeplanes, upon irradiation of a crystal in a polycrystalline specimenby an electron beam. Even for beam energies as high as 20 keV, withpenetration depths of about 2 mm in Cu(In,Ga)(S,Se)2, most of theelectrons backscattered from this volume are reabsorbed by thesurrounding material. The small information depths for EBSD haveas consequence a high spatial resolution of about 20–30 nm.

Upon electron irradiation, backscattered electrons diffracted ata certain family of lattice planes in a crystal emerge from thespecimen. The ensemble of the various diffraction bands upondiffraction at the various families of lattice planes in a crystal, theEBSD pattern, contains information not only on the symmetry ofthe crystal (and thus allows for the estimate of its phase) but also onthe crystal orientation with respect to a defined reference coordi-nate system. Therefore, by the acquisition of EBSD patterns by useof a video camera during the scanning of the electron beam across agiven specimen area, EBSD maps containing local crystal orienta-tions and symmetries are obtained. These maps allow for quantify-ing average grain sizes, local and integral textures as well as giveinformation on positions and misorientations of grain boundaries.Further details may be found in Ref. [7], and examples of EBSD mapswill be demonstrated further below in section 3.1.1.3.

3.1.1.2. Energy-dispersive X-ray spectrometry. Phase analysis bymeans of EBSD alone remains often ambiguous, owing to manyphases having similar crystal symmetries which may not be dis-tinguished by EBSD. In order to improve the quality of the analysis,elemental distributions in a given sample may be studied by meansof EDX. The combination of EBSD and EDX allows in most cases forunambiguous phase analysis [8].

It is often argued that for EDX measurements in SEM, the spatialresolution is too small for the analysis of elemental distributions inthe thin films of a Cu(In,Ga)(S,Se)2 solar-cell stack. This is, however,not true. The spatial resolution of an EDX measurement is limitedby two quantities: the acceleration voltage and the energy of the X-rays which are evaluated. The volume of excitation is defined by theacceleration voltage; the smaller this volume, the higher the spatialresolution. Also, the smaller the energy of the generated char-acteristic X-rays, the shorter is their mean-free path, and the higheragain the spatial resolution of the EDX measurement. It has beenshown [9] that spatial resolutions of down to 100 nm and belowmay be obtained in SEM–EDX measurements by reducing theacceleration voltage down to few kV.

As an example, Fig. 1 shows an elemental-distribution map,superimposed on a secondary-electron image, acquired on a cross-section specimen from a ZnO/CdS/CuInS2/Mo/glass solar cell (CIS-1,see Table 1). The map is composed of the individual maps contain-ing Zn-L, Cd-L, Cu-L and Mo-L signals. The acceleration voltage forthe acquisition of these maps was 7 kV. It is apparent that the

Fig. 1. EDX elemental distribution map superimposed on a SEM image, acquired on

a cross-section specimen from a ZnO/CdS/CuInS2/Mo/glass solar cell (CIS-1, see

Table 1). The map is composed of individual Zn-L, Cd-L, Cu-L and Mo-L

distribution maps.Fig. 2. SEM image (a), EBSD orientation distribution (b) and pattern-quality maps

(c), also with S3 grain boundaries highlighted by red lines (d) as well as a

panchromatic CL image (e) all acquired at the identical position of a ZnO/CdS/

CuGaSe2/Mo/glass cross-section specimen (CGSe-1, see Table 1). (For interpretation

of the references to color in this figure legend, the reader is referred to the web

version of this article.)

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–14621454

signals from the thin CdS layer (30 nm thick) can be clearlydistinguished from those attributed to the ZnO and CuInS2 layers.The thickness of the CdS layer appears larger in the EDX map thanits real value because of the lateral extension of the excitationvolume at 7 kV.

It is also shown by the EDX elemental-distribution map in Fig. 1that modern evaluation software suites are able to deconvolutesuccessfully overlapping X-ray peaks in the EDX spectrum, e.g., theCu-L and Zn-L, and also the S-K (not shown here) and Mo-L lines.Because of the rather strong roughness of the CuInS2 layer, theelectron beam impinging on the CdS layer excites also Zn signalsfrom ZnO. This is why the CdS layer does not appear contiguous,and also why Zn seems to be present at positions in between therough ZnO layer.

3.1.1.3. Correlation between electron backscatter diffraction and

cathodoluminescence. Apart from EDX for phase analyses, EBSDmay also be combined with other SEM techniques. Fig. 2 showscross-sectional SEM and panchromatic CL images as well as EBSDorientation-distribution and pattern-quality maps from theidentical position of a ZnO/CdS/CuGaSe2/Mo/glass solar-cell stack(sample CGSe-1, see Table 1). The local orientations in Fig. 2 b aregiven by colors, see the legend. In the pattern-quality map shown inFigs. 2d, S3 grain boundaries are highlighted by red lines (pleaserefer to Ref. [7] and to the references therein for explanationsof the S value and further details on S3 grain boundaries inCu(In,Ga)(S,Se)2 thin films).

It is important to point out that the EBSD maps in Fig. 2(b)–(d)are different representations of the identical EBSD data. Apart fromlocal orientations recorded in each pixel of the EBSD maps(Fig. 2(b)), the pattern quality can be determined by extractingprofiles across the diffraction bands in the EBSD pattern andrelating the slope of the (ideally rectangular) profile to a grayvalue. At grain boundaries, low gray values are obtained sincesuperimposing EBSD patterns from neighboring grains arerecorded, overall leading to a poor pattern quality. Since theorientations of two adjacent grains are known, their misorienta-tions (usually expressed by a rotation about a crystal axis throughan angle, transforming the point lattice of the one grain into that ofthe other) can be determined.

S3 grain boundaries in Cu(In,Ga)(S,Se)2 (Fig. 2(d)) exhibit specificmisorientations [7] and are generally of high symmetry (low densityof crystal defects). About 50 % (and above) of the grain boundaries in

Cu(In,Ga)(S,Se)2 thin films produced by co-evaporation from theelements or by an RTP are counted among this type [12]. This findingmay be explained by assuming that most of theS3 grain boundarieshave a twin constellation, where twin boundaries are special casesof stacking faults. The stacking-fault formation enthalpies inCu(In,Ga) Se2 were shown to be considerably small [10]. Thus, theformation of these planar defects represents an appropriate mechan-ism in order to reduce the strain present in the Cu(In,Ga)(S,Se)2

thin films during growth and probably also during the cooling-down stage.

The panchromatic CL image (Fig. 2(e)) exhibits an intensitydistribution which can be attributed to the microstructure repre-sented by the EBSD pattern-quality map (Fig. 2(c)). It appears thatthe CL signal intensities are higher from certain grains as comparedwith those from others. By comparison of the EBSD orientation-distribution map (Fig. 2(b)) and the CL image (Fig. 2(e)), it can beexcluded that the CL intensity is influenced by the local orientation.Recently, it was shown [11] that the CL intensity emitted fromCu(In,Ga)(S,Se)2 may be considerably influenced by the presence oflinear or planar crystal defects, probably of dislocations. Also, thedistribution of the CL signal intensity across the polycrystallineCuGaSe2 layer may be affected by slight variations in the localdoping since the radiative recombination rate depends on thecharge-carrier concentration.

It was also demonstrated [11,12] that S3 (twin) grain bound-aries in Cu(In,Ga)(S,Se)2 do not influence considerably CL inten-sities, i.e., the radiative recombination is similar as in the graininteriors, in contrast to non-S3 grain boundaries, where decreasedCL signal intensities were found. These results can be alsoconfirmed at few positions in Fig. 2(e); but overall, the signal tonoise ratio in this CL image (Fig. 2(e)) is too low for extensiveanalysis of the CL-signal changes at CuGaSe2 grain boundaries.Further analyses of grain boundaries in Cu(In,Ga)(S,Se)2 thin filmsare reported further below in section 3.2.

Generally, the CL signals are emitted upon generation ofelectron–hole pairs by the incident electron beam, diffusion ofthese charge carriers and their radiative recombination. That is, thespatial resolutions of CL images depend on the diffusion lengthsand thus on the lifetimes of the minority charge carriers. Theseresolutions can be as low as 50–100 nm and are affected by the

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–1462 1455

temperature of the specimen. The CL intensities do not only dependon the fraction of radiative recombination in the total recombina-tion of charge carriers, but also on the generation of electron–holepairs, which may be influenced by the band-gap energy ofsemiconductors or the absorption coefficients of the materialsunder investigation. When CL spectra acquired in each measuringpoint of a map are compared, changes in phases may be detected aswell as point-defect densities and their energy positions. Time-resolved CL with pulsed incident electron beams can provide accessto lifetimes of charge carriers.

Fig. 3. (a) Cross-sectional BF-TEM image of a ZnO/CdS/Cu(In,Ga)Se2/Mo/ glass solar-

cell stack (CIGSe-1, see Table 1); (b) identical image, with the stacking faults and

microtwins highlighted by red as well as the dislocations by blue lines. Figures a and

b are two BF-TEM images stitched together (cf. vertical line). (c) A DF image from the

identical Cu(In,Ga)Se2 layer shown in (a) and (b), with dislocations visible as white

lines within a Cu(In,Ga)Se2 grain.

3.1.2. Analyses of structural defects and elemental distributions by

transmission electron microscopy techniques

Due to its high spatial resolution, TEM is a powerful tool tocharacterize crystal defects in polycrystalline semiconductors [13].In particular, it is used to identify and distinguish various types ofcrystal defects such as dislocations, stacking faults and alsoantiphase or grain boundaries. Among these types, dislocationsplay an important role in semiconductors because of theirelectrically active centers in the crystal structure, generated bydangling bonds and displacement fields [14].

The characterization of dislocation networks includes thedetermination of their densities and crystal orientations. A dis-

location is defined by its line direction and its Burgers vector b!

,standing for the resultant displacement of the lattice. The dis-placement field around the dislocation changes the Bragg diffrac-tion condition, which leads to a change in the diffraction contrast.This is why dislocations appear as black lines in TEM bright-field(BF) images, as in the Cu(In,Ga)Se2 layer in Fig. 3(a), where aBF-TEM cross-section image from a ZnO/CdS/Cu(In,Ga)Se2/Mo/glass stack (CIGSe-1, see Table 1) is given. In addition, also stackingfaults and microtwins are visible (however, they cannot bedistinguished from one another in this image, only by high-resolution imaging). These structural defects are highlighted byblue and red lines in Fig. 3(b).

In contrast to BF mode, in dark-field (DF) TEM, only that areaappears bright which satisfies the Bragg condition of the beam,selected by an aperture applied in the back-focal plane of theobjective lens. Fig. 3(c) shows a DF-TEM image from an area in theCu(In,Ga)Se2 layer (identical specimen as for Fig. 3(a) and (b)), withdislocations visible as narrow bright lines on dark background. TheDF conditions were selected in a way that only the surroundeddisplacement field satisfies the Bragg condition in contrast to theorientation of the grain. Furthermore, one can use special DFanalyses to determine the dislocation parameters such as theburgers vector and the line direction as well as the dissociationwidth of partial dislocations [13]. TEM is thus able to provide adetailed structural characterization of one- and two-dimensionallattice defects.

TEM analyses of various Cu(In,Ga)(S,Se)2 layers deposited by co-evaporation from the elements or by means of an RTP show aconsiderably high relative frequency of twin boundaries andstacking faults, as already found by means of EBSD (Section3.1.1.3). It is not yet clear what the formation mechanisms ofstacking faults and microtwins are, and whether they are morepreferentially formed than the dislocation networks shown inFig. 3. This issue is currently under investigation.

In addition to microstructural studies, compositional analysesmay be performed using EDX or also by energy electron energy-lossspectroscopy (EELS) in TEM. While EDX has been applied to a largeextent in investigations on Cu(In,Ga)(S,Se)2 thin-films and corre-sponding solar cells, only few reports can be found in the literatureon EELS analyses [15–17]. This is mainly due to the fact that thecorresponding loss edges for Cu, Ga and Se, used for elementalimaging, are positioned at rather high energy losses ranging from

about 930 to 1500 eV, leading to low signal intensities and thus tolow signal-to-noise ratios and overall unfavorable conditions withrespect to EDX.

Fig. 5. EDX elemental distribution profiles, extracted from corresponding EDX maps

across an CuInS2/CdS interface (CIGSe-12, see Table 1). The net counts were

normalized to the maximum intensity of the Cd-L signal, for easier comparison.

The net counts are given as solid symbols, the solid lines are smoothed data given as

guide for the eye. Apparently, the In-L decreases more gradually than the Cu-K signal,

indicating Cu diffusion CuInS2 from into CdS. The Cd-L signal is not zero in the CuInS2,

but it remains ambiguous because of possible superposition with the In-L signal.

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–14621456

Due to the very small thicknesses of the TEM specimens studied(50–100 nm), the extension of the generation volume of theelectron beam for an EDX measurement in TEM is one or twoorders of magnitude smaller compared with EDX in SEM (Section3.1.1.2), which leads to a much higher spatial resolution of fewnanometers (nowadays, by use of aberration-corrected micro-scopes, values of down to 0.4–0.6 nm have been realized [18]).Disadvantages of the small specimen thicknesses are the generallypoor measurement statistics and the sensitivity to thicknessvariations.

Fig. 4 is composed of the EDX elemental distribution maps usingZn-K, Cd-L, Ga-K and Mo-K signals, superposing a scanning TEM(STEM) image, from a ZnO/CdS/Cu(In,Ga)Se2/Mo/glass solar-cellstack (CIGSe-1, see Table 1). The position of the CdS layer is muchbetter visible than from the maps in Fig. 1. Also, the Ga signalexhibits an intensity gradient across the Cu(In,Ga)Se2 layer, withlocal maxima close to the interfaces to the Mo and CdS layers. Suchelemental distribution maps are essential for the research anddevelopment of Cu(In,Ga)Se2 solar cells, giving information on localcompositions and interdiffusions down to few nanometers (inuncorrected microscopes).

A further example is shown in Fig. 5. The EDX linescans given inthere show Cu, In, Ga and Cd distribution profiles, extracted fromcorresponding Cu-K, In-L and Cd-L distribution maps acquired on aCuInS2/CdS interface. It is clearly visible that the In signal decreasesmore gradually than the Cu signal, where the drop-off of the Cusignal is shifted towards the CdS signal, suggesting Cu diffusionfrom CuInS2 into CdS. Also, the Cu signal is still enhanced where theIn signal intensity is already decreased to zero. Indeed, such adiffusion has been found by various authors [19,20] for CBD-CdSbuffer layers grown on Cu(In,Ga)Se2.

3.2. Grain-boundary physics in Cu(In,Ga)(S,Se)2 thin films

Although there is a high density of grain boundaries in poly-crystalline Cu(In,Ga)Se2 layers, solar-conversion efficiencies of upto more than 20 % have been achieved [21]. There are variousmodels, simulations and experimental results treating the influ-ence of grain boundaries and explaining why they may not have adetrimental effect on charge-carrier collection and thus on thedevice performance [22–34]. Some of these models and results[22,28,29] even suggest that grain boundaries have a positive effectoriginating from charged or neutral hole or electron barriers

Fig. 4. EDX elemental distribution maps composed by use of Zn-K, Cd-L, Ga-K and

Mo-K signals, superposing a STEM image, acquired on a ZnO/CdS/Cu(In,Ga)Se2/Mo/

glass solar-cell stack (CIGSe-1, see Table 1).

present in their proximity, which would impede recombinationof electrons and holes.

The following subsections describe approaches for exploring theeffects of grain boundaries in Cu(In,Ga)Se2 layers on charge-carriercollection. It is important to note that the corresponding electronmicroscopy results are rather related to grain-boundary propertiesin the volume of these Cu(In,Ga)Se2 layers, in contrast to scanningprobe techniques analyzing the ones at their surfaces. While at thegrain boundaries on the surface, charge accumulation is probable[24,26], this is not necessarily the case for grain boundaries in thevolume [35], as we will also show further below.

3.2.1. Combination of electron-beam-induced current and electron

backscatter diffraction measurements

In this section, electron-beam induced current (EBIC) measure-ments and EBSD in the cross-section and backside configurationsare presented as capable techniques in order to gain insightinto charge-carrier collection properties of grain boundaries inCu(In,Ga)Se2 layers of thin-film solar cells. Due to the relativelysmall extension of the generation volume of the electron beam ascompared with generation by light, EBIC is an adequate tool tostudy charge-carrier collection spatially resolved, while EBSDserves to localize and classify grain boundaries [7]. It can be shown(see below) that effects of single grain boundaries in standardCu(In,Ga)Se2 absorber layers with an average grain size and anelectron diffusion length in the mm range can be identified bymeans of EBIC when using low electron beam energies, ensuring asmall extension of the generation volume and Green penetrationdepth [36] (e.g., about 200 nm for 5 keV).

Problems in the interpretation of EBIC results may arise due tothe fact that the data is only accessible in two dimensions, whereascharge-carrier collection is a process taking place in three dimen-sions. Also, there are additional surfaces present, where enhancedrecombination or charge accumulation is possible.

3.2.1.1. Introduction to EBIC. Fig. 6 shows schematically the setup ofan EBIC experiment. The solar-cell sample is irradiated by a focusedelectron beam, and charge carriers are generated in a defined regionaround the position of irradiation. Empirical expressions for the

Fig. 6. Schematic representation of EBIC setups for backside (a) and cross-section

(c) configurations. The depth dependent (b) and lateral (d) generation profiles

in a Cu(In,Ga)Se2 layer as well as the corresponding, calculated collection functions

fc in Cu(In,Ga)Se2 solar cells are also shown. fc is assumed to be 0 in the n-type layers,

1 in the SCR and is distributed according to Eq. (3) in the QNR of the Cu(In,Ga)

Se2 layer.

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–1462 1457

description of the lateral and depth-dependent generation profilesare given in Refs. [36,37]. The smaller the beam energy Eb, thenarrower is the lateral generation profile (see Fig. 6). Thus, foridentifying effects of individual grain boundaries, it is necessary touse low beam energies. However, there is a trade off between spatialresolution and the influence of surface effects, which has to be takeninto account. Therefore, beam energies between about 4 and 10 kVare recommended for average grain sizes larger than 500 nm.

Most relevant in order to quantify charge-carrier collection

processes is the collection function fcð x!Þ, which stands for the

probability of a charge carrier generated at position x!

to becollected. The measured EBIC signal can be expressed as a

convolution of the collection function fcð x!Þ and the generation

function gð x!

, a!Þ of the electron beam ( a

!is the impinging position

of the electron beam) [38]

Ið a!Þ¼

Zfcð x!Þgð x!

, a!Þd x!

ð1Þ

A solution for the collection function fcð x!Þ for the quasi-neutral

region (QNR) of the absorber layer under low injection conditionscan be gained from the following differential equation, which wasderived from the continuity equation via a reciprocity theorem [38]

DDfcð x!Þþm E!!fcð x!Þ�

fcð x!Þ

t¼ 0 ð2Þ

where D is the minority charge carrier (i.e., electron) diffusionconstant, m the electron mobility, E

!the electrical field and t the

electron lifetime. If translation invariance in the directions parallelto the p–n junction is assumed, a one-dimensional equation can beused. The effects of grain boundaries are not covered in thisapproach; for simplification, it is necessary to neglect them at thispoint. As a boundary condition, it is assumed that the collectionprobability at the edge of the space-charge region (SCR) is 1(fc(xSCR)¼1) and that fc(x)-0 for x-N for an infinite semiconduc-tor layer. In this case, the solution for the collection function fcð x

!Þ is

a simple exponential function fc(x)¼exp(�x/L), where L is theelectron diffusion length. Assuming a finite semiconductor limitedby a back contact at position xBC, the second boundary conditionchanges to dfc=dxðxBCÞ ¼ SBC=DfcðxBCÞ, where SBC is the electronrecombination velocity at the back contact. A solution for fc (x) is

fcðxÞ ¼1=Lcoshððx�xBCÞ=LÞ�SBC=Dsinhðx�xBCÞ=LÞ

SBC=DsinhððxBC�xSCRÞ=LÞþð1=LÞcoshððxBC�xSCRÞ=LÞð3Þ

In this study, we present EBIC measurements performed on thebackside of a Cu(In,Ga)Se2/CdS/ZnO stack lifted off from the Mo/glass substrate (Fig. 6(a)), and on a polished cross-section of theZnO/CdS/Cu(In,Ga)Se2/Mo/glass solar-cell stack. In Fig. 6(b) and (d)(red curves), exemplary collection functions of a Cu(In,Ga)Se2

thin-film solar cell are shown. fc(x) is assumed to be 0 in thewindow and buffer layers, and 1 in the SCR. The diffusion lengthand the recombination velocity at the back contact are L¼500 nmand SBC¼1�105 cm/s.

In former studies, EBIC results on Cu(In,Ga)Se2 solar cells with CdSbuffer layers were found to deviate from the theoretical expectationsgiven above. This deviation was explained [39] by assuming thatelectrical defects and in consequence charges accumulate at the CdS/Cu(In,Ga)Se2 interface, causing a configuration where the collectionfunction is not independent of the generation function. This has to betaken into account when evaluating EBIC data.

3.2.1.2. EBIC signals at grain boundaries. In case of enhancedrecombination at a grain boundary, a local minimum in the EBICsignal across the grain boundary is expected, because charge

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–14621458

carriers generated at or close to the grain boundary exhibit ashorter effective lifetime and are more likely to recombine thancharge carriers generated in the interior of the Cu(In,Ga)Se2 grain.The equations given in Refs. [40,41] provide an expression for areduced, effective diffusion length measured in the proximity of thecorresponding grain boundary. The shape of an EBIC linescanperpendicular to a grain boundary can be calculated from theseequations, and by comparing experimental data with the simulatedcurves, an estimate of the recombination velocity at the grainboundary, Sgb, is possible. Fig. 7 shows simulated EBIC profilesperpendicular to a grain boundary in the cross-section config-uration assuming different recombination velocities (for details seeRef. [27]). These simulations also show that the spatial resolution ofEBIC is sufficient to resolve effects of individual grain boundarieswhen using low beam energies of around 5 keV and for averagegrain sizes in the range of 0.5 mm or larger. When using higherbeam energies (e.g. 15 keV) and in case average grain sizes aresignificantly smaller, the effects of various grain boundariessuperimpose and the EBIC signal smears out.

In case of a neutral hole barrier at a grain boundary [22], reducedrecombination of electrons and therefore enhanced charge-carriercollection is expected. The widths of the resulting EBIC maxima atsuch grain boundaries depends on the electron-beam energy Eb

applied and on the electron diffusion-length L in Cu(In,Ga)Se2 as incase of enhanced recombination.

Fig. 7. Simulated normalized EBIC profiles perpendicular to a grain boundary of the

Cu(In,Ga)Se2 layer. The current values were extracted from simulated EBIC line

profiles perpendicular to the pn-junction in a distance of 200 nm to the edge of the

SCR in the QNR of the Cu(In,Ga)Se2 layer. The corresponding effective diffusion

lengths for the different recombination velocities Sgb, the different distances to the

grain boundary and the different electron-beam energies were calculated from the

equations given in Refs. [40,41].

If a grain boundary is positively charged, a SCR develops, whichmay assist in charge-carrier collection by separating generatedelectrons and holes [28]. Depending on the doping density of theCu(In,Ga)Se2 layer, this separation may lead to either a maximumin the EBIC linescan in the proximity of the grain boundary (broader

Fig. 8. SEM image (a), EBSD orientation-distribution map (b) with local orientations

given by colors (see legend), and EBIC image acquired at 8 keV (c) on the identical

sample position of a polished cross-section of a Cu(In,Ga)Se2 thin-film solar cell

(CIGSe-2, see Table 1). The white arrows indicate the position of the extracted EBIC

profiles shown in (d). When comparing the experimental data with the simulations

(d), a recombination velocity of 3�104 cm/s can be estimated for this grain

boundary. For details, see Ref. [27].

Fig. 9. SEM (a) and EBIC images (b) as well as a EBSD orientation-distribution map

(c) from the identical sample position on the backside of Cu(In,Ga)Se2/CdS/ZnO stack

(prepared from CIGSe-3, see Table 1). The white arrows in (b) and (c) indicate the

position of the extracted EBIC profile (d) perpendicular to a grain boundary at EB ¼

8 keV. IEBIC/IB is the ratio of measured EBIC and electron-beam currents.

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–1462 1459

than in the case of the neutral hole barrier) or a uniform EBIC signal(see also Ref. [27]).

The results shown in the following are not considered repre-sentative for all grain boundaries in Cu(In,Ga)Se2 layers; theintention here is rather to demonstrate how to apply EBSD andEBIC on identical sample positions.

3.2.1.3. Cross-sectional electron-beam-induced current measure-

ments. Fig. 8 shows SEM and EBIC images as well as an EBSDorientation-distribution map from the identical position of thepolished cross-section of a ZnO/CdS/Cu(In,Ga)Se2/Mo/glass solarcell (CIGSe-2, see Table 1). The white arrow indicates the position ofthe extracted EBIC profile perpendicular to a grain boundary, whichis shown in Fig. 8(d). The current values of these profiles weredetermined by first extracting EBIC profiles perpendicular to thep–n junction at different distances to the grain boundary. Then, theposition of the edge of the SCR of each of these EBIC profiles wasdefined and the current value of this position was set to 1. For theresulting normalized EBIC profile perpendicular to the grainboundary and parallel to the p–n junction (Fig. 8d), the currentvalues were extracted at a distance of 200 nm from the edge of theSCR in the QNR of the Cu(In,Ga)Se2 layer and normalized to theplateau on the right hand side, which was defined to be 1. Forfurther details, see Ref. [27]. This procedure has to be carried outsince the model explained above takes into account only differ-ences in the EBIC profile perpendicular to the p–n junction inthe QNR.

The resulting EBIC profile exhibits a local minimum at theposition of this grain boundary. A minimum was also found forhigher beam energies Eb (see Ref. [27]). When comparing thedecrease of the EBIC signal intensity with the simulated curves [27],a recombination velocity of 3�104 cm/s was estimated for thisgrain boundary. This is a rather low value.

3.2.1.4. Backside electron-beam-induced current measurements. Bylifting off the glass substrate and the Mo layer the Cu(In.Ga)Se2

solar cell is accessible from the backside of the absorber layer.Again, EBSD and EBIC measurements were performed on theidentical sample position (see Fig. 9). The EBIC distribution profilegiven in Fig. 9(d) is an average of 30 individual profiles and shows aminimum (decrease of 5 rel.%) at the grain-boundary position.Similarly as for the analyses in cross-section configuration (Fig. 8),significant changes in the EBIC signal have only been found at non-S3 grain boundaries. However, the simple, one-dimensional ana-lytical model applied for the EBIC profiles in the cross-sectionconfiguration was unable to describe satisfactorily the profilesextracted from the backside EBIC images.

It may be noted that EBSD and EBIC results also from otherbackside Cu(In,Ga)Se2/CdS/ZnO samples with varying Ga and Cuconcentrations [42] showed a reduced short-circuit current mea-sured upon electron-beam irradiation on non-S3 grain boundaries.

In general, it was found that the EBIC signal does not changesignificantly at the positions of S3 (twin) grain boundaries,whereas there are different effects observed at randomly orientedgrain boundaries of Cu(In,Ga)Se2 absorber layers. The resultsshown above are only examples and serve to demonstrate theexperimental approach. For a general statement about the influ-ence of grain boundaries more data needs to be collected. Addi-tionally, two- and three-dimensional numerical device modellingis needed and will be applied for future studies.

3.2.2. Grain boundaries analysed by inline electron holography

Inline electron holography in TEM [3,43] is a valuable tool toinvestigate the electromagnetic potentials at interfaces on a subnanometer scale. In the present section, we apply this technique for

Fig. 10. Gray-value image of the reconstructed phase differenceDj(x,y), acquired at

a random (non-S3) grain boundary (GB) within a CuInS2 absorber (CIS-2, see

Table 1). Individual profiles were extracted across grain boundaries, from areas

similar to that depicted by the black frame.

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–14621460

studying the behavior of the Coulomb potential at grain boundariesin polycrystalline Cu(In,Ga)Se2 and CuInS2 solar cell absorbers.

In TEM, the electron beam can be described by a plane wavefunction that impinges on a specimen. The interaction with thespecimen then alters the amplitude and the phase of the wavefunction. The phase of the wave function at the exit plane of thespecimen contains information about the electromagnetic poten-tial within the specimen. However, this information cannot beutilized in conventional transmission electron microscopy since forthis microscope mode, only the intensity distribution is recorded,which is given by the squared modulus of the image wave functionthat arises from the transfer of the exit plane wave function by themicroscopic system.

Thus, all phase information is lost. However, the phase j(x,y) asfunction of the local position (x,y) within the exit plane may beretrieved by using holographic methods. We applied inline electronholography (see Ref. [44] for a comparison with off-axis electronholography). I.e., we acquired a through-focal series of images,from which the relative phase distribution Dj(x,y) with respect toan unknown reference phase value was reconstructed [3].

In the following, we will apply the absorbing phase objectapproximation, in which multiple elastic scattering and also inelas-tic scattering out of the elastic channel are treated to all orders, butthe curvature of the Ewald sphere is effectively neglected. Thisapproximation is therefore suitable for zero-loss, energy-filteredimages at medium resolution as were acquired for this work. Thephase of the impinging electron plane wave is affected by theinteraction with the specimen, because the wavelength of the wavefunction depends on the local electrostatic and vector potentialwithin the specimen. Assuming that there is no magnetic flux withinthe specimen, the wavelength l is given by [45]

l¼hffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

2m0eðUþVÞð1þðeðUþVÞÞ=ð2m0c2ÞÞp ,

where h is Planck’s constant, m0 and e are the electron rest mass andcharge, U the acceleration voltage of the transmission electronmicroscope and V the electrostatic potential which the electronsfrom the incident beam experience when they travel through thespecimen. It contains contributions from scattering atoms (or ions)and free charge carriers within the sample and from additionalelectric fields. At the exit plane of the TEM specimen, the change inphase j(x,y) due to the electrostatic potential V(x,y,z) within thespecimen, which is assumed to be small compared with the kineticenergy of the beam electrons, is then given by [45]

jðx,yÞ ¼ sZ

Vðx,y,zÞdz¼ sðUÞtðx,yÞVeðx,yÞ, ð4Þ

where s(U) is the interaction constant, Ve(x,y) the electrostaticpotential averaged along the z-axis perpendicular to the specimensurface and t(x,y) the local specimen thickness. Note that Ve(x,y) isthe electrostatic potential and not the potential energy. For neutralsolids without magnetic flux, Ve(x,y) is positive, i.e., the electrons seean attractive potential within the solid. By means of inline electronholography, we cannot determine the absolute phase shift inducedby interaction with the specimen but only relative phase shiftsDj(x,y) originating from relative potential differences DVe(x,y) atvarious specimen positions.

For the acquisition of the through-focal series, we tilted thespecimen, until the planes of the grain boundaries were parallel tothe incident electron beam. We reconstructed gray-value images ofthe relative phase shiftDj(x,y) at the grain boundaries by applyingan algorithm developed by Koch [4]. Fig. 10 exemplarily showssuch a gray-value image of the reconstructed phase Dj(x,y) at anon-twin (‘‘random’’, probably also non-S3) grain boundary withina polycrystalline CuInS2 absorber.

While the grain interiors exhibit a negligible phase shifts, thephase shifts at the grain boundaries are significantly larger. Fromgray-value images of the phase difference Dj(x,y) we extractedprofiles across grain boundaries for Cu(In,Ga)Se2 and CuInS2 solar-cell absorbers (CIGSe-4 and CIS-2, see Table 1). The black frame inFig. 10 denotes the region within which about 50 profiles wereaveraged parallel. From these profiles, we calculated the relativeelectrostatic potential DVe(x,y) by applying Eq. (4).

For the measurement of the specimen thicknesses t(x,y), weapplied energy-filtered TEM. We determined maps of the specimenthickness t(x,y) in units of the inelastic mean free path lmean ofelectrons within the specimen by taking the natural logarithm ofthe ratio between an unfiltered and a zero-loss filtered image [46]t(x,y)/lmean ¼ ln(It(x,y)/I0(x,y)). It(x,y) is the intensity distributionof an unfiltered image, and I0(x,y) the intensity distribution of azero-loss filtered image, to which only unscattered or elasticallyscattered electrons contribute. lmean can be estimated by use of analgorithm given in Ref. [46].

Fig. 11 depicts the resulting electrostatic profiles (averaged)across a twin (S3) and a random (probably non-S3) grain boundaryin a Cu(In,Ga)Se2 layer as well as a random grain boundary within aCuInS2 thin film (CIGSe-4 and CIS-2, see Table 1). The potentialwells at twin boundaries are often indistinguishable from the noisein electrostatic potential profiles (about 100 mV) and were found toexhibit depths of about 200 mV at most. This indicates, correspond-ingly to the EBIC results, not any substantial changes of theelectrostatic potential at this type of grain boundary.

We measured only potential wells (and not any barriers) at grainboundaries in various Cu(In,Ga)Se2 and CuInS2 thin films incompleted solar cells. While the full width at half minimum(FWHM) of all potential wells was 1–2 nm, the potential-welldepths ranged from 200 mV for the twin grain boundary inCu(In,Ga)Se2 to 1.5 V for random grain boundaries in CuInS2. Ingeneral, from studies on various other Cu(In,Ga)Se2 thin films,values between 1 and 3 V were obtained.

Fig. 11. Extracted electrostatic potential profiles across grain boundaries in

Cu(In,Ga)Se2 and CuInS2 absorbers (CIGSe-4 and CIS-2, see Table 1). These profiles

are averages of 50 individual profiles extracted from areas similar to the black frame

given in Fig. 10.

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–1462 1461

There are several possible reasons for the presence of potentialwells at grain boundaries. First, a negative charge at the grain-boundary core may be considered, which is screened by positivecharges in the grain interiors. For typical doping concentrations of1016–1017 cm�3 in Cu(In,Ga)Se2 and CuInS2, the correspondingDebye lengths can be calculated to be between 12 and 44 nm atroom temperature. However, the FWHM of the potential wells ismuch smaller than these values. With the reasonable assumptionthat the electron beam does not generate a free-charge carrierdensity exceeding that of the doping (the electron-beam intensitywas reduced accordingly), the potential-well shape cannot beexplained only by the screening of negative charges located atthe grain-boundary cores.

As second possibility, a space-charge region screening a positivecharge at the grain boundary core is possible, but would result in apotential barrier and thus, does not agree with our results.However, we cannot exclude from our measurements that apotential barrier smaller than about 100 meV is present, whichmay be caused by a positive charge at the grain boundary. The noiseof the electrostatic potential signals is in about the same range.

It seems therefore likely that the main contribution to thepotential wells obtained by inline electron holography are notcharge accumulations but rather changes of crystal structure andcomposition at Cu(In,Ga)(S,Se)2 grain boundaries. This issue iscurrently under investigation in order to provide evidence forthese compositional changes. It is also to be clarified what thelink of these reconstructed grain boundaries are to the reducedshort-circuit current measured by means of EBIC.

4. Conclusions

In the present work, various SEM and TEM techniques arereported which may be applied for the analysis of Cu(In,Ga)(S,Se)2

solar cells. It was shown that each technique gives importantinformation on microstructure, composition as well as on electricaland optoelectronic properties, but that combinations of electron-microscopy methods improve considerably the scientific quality ofthe results. EBSD may be combined with EDX for unambiguousphase analysis. In TEM, BF and DF imaging reveal linear or planardefects, and compositional gradients in Cu(In,Ga)Se2 layers as wellas interdiffusion between Cu(In,Ga)Se2 and buffer layers may bestudied by means of EDX.

At grain boundaries, EBSD maps together with EBIC and CLmeasurements performed on identical positions give informationon the corresponding charge-carrier collection and radiativerecombination. For non-S3 grain boundaries in Cu(In,Ga)(S,Se)2

thin films, reduced EBIC and CL signals were found, with estimatedrecombination velocities of few 104 cm/s. We would like to notethat in few samples deviations from these findings were alsoobtained. In contrast, not any considerable electronic activity wasdetected at S3 (twin) boundaries. This result was confirmed byinline electron holography, by which only at non-S3 grain bound-aries in Cu(In,Ga)(S,Se)2 thin films, Coulomb-potential wells withwidths of 1–2 nm and depths of 1–3 V were measured. Overall,the results presented here indicate an atomic reconstructionvia neutral defects or defect complexes at grain boundaries inCu(In,Ga)(S,Se)2, which only slightly reduce the short-circuit current.

Acknowledgements

The authors would like to thank T. Unold, H.-W. Schock, and C.Boit for fruitful discussions and continuous support. Special thanksare due to B. Bunn, C. Kelch, M. Kirsch, and T. Munchenberg for helpin sample preparation, as well as to J. Bundesmann, H. Stapel and M.Wollgarten for scientific and technical assistance with the electronmicroscopes. The authors would like to acknowledge funding fromthe German ministry for environment (BMU) under contract#0327559H (German-Israeli project ProGraB) as well as from theGerman ministry for Education and Research (BMBF) under con-tract #033F0359C (project GRACIS).

References

[1] C.A. Kaufmann, R. Caballero, T. Unold, R. Hesse, R. Klenk, S. Schorr,M. Nichterwitz, H.-W. Schock, Depth profiling of Cu(In,Ga)Se2 thin filmsgrown at low temperatures, Sol. Energy Mat. Sol. Cells 93 (2009) 859–863.

[2] K. Siemer, J. Klaer, I. Luck, J. Bruns, R. Klenk, D. Braunig, Efficient CuInS2 solarcells from a rapid thermal process (RTP), Sol. Energy Mat. Sol. Cells 67 (2001)159–166.

[3] S. Bhattacharyya, C.T. Koch, M. Ruhle, Projected potential profiles acrossinterfaces obtained by reconstructing the exit face wave function from throughfocal series, Ultramicroscopy 106 (2006) 525–538.

[4] C.T. Koch, A flux-preserving non-linear inline holography reconstructionalgorithm for partially coherent electrons, Ultramicroscopy 108 (2008)141–150.

[5] A.J. Schwartz, M. Kumar, B.L. Adams, D.P. Field (Eds.), Electron backscatterdiffraction in materials science, Springer, New York, , 2009.

[6] D. Abou-Ras, R. Caballero, J. Kavalakkatt, M. Nichterwitz, T. Unold, H.W. Schock,S. Bucheler, A.N. Tiwari, Electron backscatter diffraction: exploring the micro-structure in Cu(In,Ga)(S,Se)2 and CdTe thin-film solar cells, in: ConferenceRecords of the 35th IEEE Photovoltaics Specialists Conference, Honolulu,Hawai’i, June 20–25, 2010, IEEE, Piscataway, NJ, pp. 418–423.

[7] D. Abou-Ras, S. Schorr, H.W. Schock, Grain-size distributions and grainboundaries of chalcopyrite-type thin films, J. Appl. Cryst. 40 (2007) 841–848.

[8] D. Abou-Ras, M. Nichterwitz, R. Caballero, C.A. Kaufmann, T. Unold, S. Schorr, R.Scheer, J. Klaer, H.-W. Schock, (Enhanced) insight in the microstructure andcomposition of chalcopyrite-type thin film solar cells, in: Proceedings of the22nd European Photovoltaic Solar Energy Conference and Exhibition, Milan,Italy, September 3–7, 2007, WIP Munich, 2007, pp. 1911–1914.

[9] I. Barkshire, P. Karduck, W.P. Rehbach, S. Richter, High-spatial-resolution low-energy electron beam X-ray microanalysis, Mikrochim. Acta 132 (2000)113–128.

[10] N.I. Medvedeva, E.V. Shalaeva, M.V. Kuznetsov, M.V. Yakushev, First-principlesstudy of deformation behavior and structural defects in CuInSe2 and Cu(In,Ga)Se2,Phys. Rev. B 73 (2006) 035207-1–035207-6.

[11] D. Abou-Ras, U. Jahn, M. Nichterwitz, T. Unold, J. Klaer, H.-W. Schock, Combinedelectron backscatter diffraction and cathodoluminescence measurements onCuInS2/Mo/glass stacks and CuInS2 thin-film solar cells, J. Appl. Phys. 107(2010) 014311-1–014311-8.

[12] D. Abou-Ras, C.T. Koch, V. Kustner, P.A. van Aken, U. Jahn, M.A. Contreras,R. Caballero, C.A. Kaufmann, R. Scheer, T. Unold, H.-W. Schock, Grain-boundarytypes in chalcopyrite-type thin films and their correlations with film textureand electrical properties, Thin Solid Films 517 (2009) 2545–2549.

[13] C.J. Kiely, R.C. Pond, G. Kenshole, A. Rockett, A. TEM, study of the crystal-lography and defect structures of single crystal and polycrystalline copperindium diselenide, Philos. Mag. A 63 (1991) 1249–1273.

D. Abou-Ras et al. / Solar Energy Materials & Solar Cells 95 (2011) 1452–14621462

[14] D.C. Holt, B.G. Yacobi, Extended defects in semiconductors, chapter 5: Theelectrical, optical and device effects of dislocations and grain boundaries,Cambridge University Press, 2007.

[15] R. Saez-Araoz, D. Abou-Ras, T.P. Niesen, A. Neisser, K. Wilchelmi, M.Ch.Lux-Steiner, A. Ennaoui, In situ monitoring the growth of thin-film ZnS/Zn(S,O)bilayer on Cu-chalcopyrite for high performance thin film solar cells, Thin SolidFilms 517 (2009) 2300–2304.

[16] T. Torndahl, E. Coronel, A. Hultqvist, C. Platzer-Bjorkman, K. Leifer, M. Edoff, Theeffect of Zn1-xMgxO buffer layer deposition temperature on Cu(In,Ga)Se2 solarcells: a study of the buffer/absorber interface, Prog. Photovolt.: Res. Appl. 17(2009) 115–125.

[17] C. Camus, D. Abou-Ras, N.A. Allsop, S.E. Gledhill, T. Kohler, J. Rappich,I. Lauermann, M.C. Lux-Steiner, C.-H. Fischer, Formation of CuInS2–carbonmultilayers in the spray ILGAR process, Phys. Status Solidi A 207 (1) (2010)129–131.

[18] M. Watanabe, D.W. Ackland, A. Burrows, C.J. Kiely, D.B. Williams, O.L. Krivanek,N. Dellby, M.F. Murfitt, Z. Szilagyi, Improvements in the X-ray analyticalcapabilities of a scanning transmission electron microscope by spherical-aberration correction, Microsc. Microanal. 12 (2006) 515–526.

[19] T. Nakada, A. Kunioka, Direct evidence of Cd diffusion into Cu(In,Ga)Se2 thinfilms during chemical-bath deposition process of CdS films, Appl. Phys. Lett. 74(17) (1999) 2444–2446.

[20] D. Abou-Ras, G. Kostorz, A. Romeo, D. Rudmann, A.N. Tiwari, Structural andchemical investigations of CBD- and PVD-CdS buffer layers and interfaces inCu(In,Ga)Se2-based thin film solar cells, Thin Solid Films 480-481 (2005)118–123.

[21] M. Powalla, P. Jackson, E. Lotter, D. Hariskos, S. Paetel, R. Wurz, R. Menner,W. Wischmann, unpublished.

[22] C. Persson, A. Zunger, Anomalous grain boundary physics in polycrystallineCuInSe2: the existence of a hole barrier, Phys. Rev. Lett. 91 (26) (2003)266401-1–266401-4.

[23] J.Y.W. Seto, The electrical properties of polycrystalline silicon films, J. Appl.Phys. 46 (12) (1975) 5247–5254.

[24] S. Sadewasser, Th. Glatzel, S. Schuler, S. Nishiwaki, R. Kaigawa, M. Ch., Lux-Steiner, Kelvin probe force microscopy for the nano scale characterization ofchalcopyrite solar cell materials and devices, Thin Solid Films 431-432 (2003)257–261.

[25] D. Fuertes Marron, S. Sadewasser, A. Meeder, T. Glatzel, M.Ch. Lux-Steiner,Electrical activity at grain boundaries of Cu(In,Ga)Se2 thin films, Phys. Rev. B 71(2005) 033306-1-4.

[26] C.S. Jiang, R. Noufi, J.A. AbuShama, K. Ramanathan, H.R. Moutinho, J. Pankow,M.M. Al-Jassim, Local built-in potential on grain boundary of Cu(In,Ga)Se2 thinfilms, Appl. Phys. Lett. 84 (18) (2004) 3477–3479.

[27] M. Nichterwitz, D. Abou-Ras, K. Sakurai, J. Bundesmann, T. Unold, R. Scheer,H.W. Schock, Influence of grain boundaries on current collection inCu(In,Ga)Se2 thin-film solar cells, Thin Solid Films 517 (2009) 2554–2557.

[28] D. Azulay, O. Millo, I. Balberg, H.W. Schock, I. Visoly- Fisher, D. Cahen, Currentroutes in polycrystalline CuInSe2 and Cu(In,Ga)Se2 films, Sol. Energy Mater. Sol.Cells 91 (2007) 85–90.

[29] M. Kawamura, T. Yamada, N. Suyama, A. Yamada, M. Konagai, Grain boundaryevaluation of Cu(In1-xGax)Se2 solar cells, Jpn. J. Appl. Phys. 49 (2010)062301-1-3.

[30] K. Taretto, U. Rau, J.H. Werner, Numerical simulation of grain boundary effectsin Cu(In,Ga)Se2 thin-film solar cells, Thin Solid Films 480-481 (2005) 8–12.

[31] K. Taretto, U. Rau, Can grain boundaries improve the performance ofCu(In,Ga)Se2 solar cells?, in: T. Gessert, S. Marsillac, T. Wada, K. Durose,C. Heske (Eds.), Thin-Film Compound Semiconductor Photovoltaics 2007,Warrendale, PA, 2007, p. Y09-01-1-6 Mater. Res. Soc. Symp. Proc. 1012.

[32] M. Gloeckler, W.K. Metzger, J.R. Sites, Simulation of polycrystallineCu(In,Ga)Se2 solar Cells in two dimensions, in: W. Shafarman, T. Gessert,S. Niki, S. Siebentritt (Eds.), Thin-Film Compound Semiconductor Photovol-taics, F14, Warrendale, PA, 2005, p. 1-1-6 Mat. Res. Soc. Symp. Proc. 865.

[33] M. Gloeckler, J.R. Sites, W.K. Metzger, Grain-boundary recombination inCu(In,Ga)Se2 solar cells, J. Appl. Phys. 98 (2005) 113704-1–113704-10.

[34] W.K. Metzger, M. Gloeckler, The impact of charged grain boundaries onthin-film solar cells and characterization, J. Appl. Phys. 98 (2005)063701-1–063701-10.

[35] S. Sadewasser, D. Abou-Ras, D. Azulay, R. Baier, I. Balberg, D. Cahen, S. Cohen,K. Gartsman, G. Karuppiah, J. Kavalakkatt, W. Li, O. Millo, Th. Rissom,Y. Rosenwaks, H.-W. Schock, A. Schwarzman, T. Unold, Nanometer-scaleelectronic and microstructural properties of grain boundaries in Cu(In,Ga)Se2,Thin Solid Films, in preparation.

[36] T.E. Everhart, P.A. Hoff, Determination of kilovolt electron energy dissipationvs. penetration distance in solid materials, J. Appl. Phys. 42 (13) (1971)5837–5846.

[37] J. Rechid, A. Kampmann, R. Reinecke-Koch, Characterising superstrate CIS solarcells with electron beam induced current, Thin Solid Films 361–362 (2000)198–202.

[38] C. Donolato, An alternative proof of the generalized reciprocity theorem forcharge collection, J. Appl. Phys. 66 (9) (1989) 4524–4526.

[39] R. Kniese, M. Powalla, U. Rau, Characterization of the CdS/Cu(In,Ga)Se2

interface by electron beam induced currents, Thin Solid Films 515 (2007)6163–6167.

[40] L. Jastrzebski, J. Lagowski, H.C. Gatos, Application of scanning electronmicroscopy to determination of surface recombination velocity: GaAs, Appl.Phys. Lett. 27 (10) (1975) 537–539.

[41] Donolato, Evaluation of diffusion lengths and surface recombination velocitiesfrom electron beam induced current scans, Appl. Phys. Lett. 43 (1) (1983)120–122.

[42] J. Kavalakkatt, D. Abou-Ras, R. Caballero, M. Nichterwitz, T. Rissom, T. Unold,H.-W. Schock, unpublished.

[43] R.E. Dunin-Burkowski, The development of Fresnel contrast analysis, and theinterpretation of mean inner potential profiles at interfaces, Ultramicroscopy83 (2000) 193–216.

[44] C.T. Koch, A. Lubk, Off-axis and inline electron holography: A quantitativecomparison, Ultramicroscopy 110 (2010) 460–471.

[45] L. Reimer, Transmission Electron Microscopy, 4th ed., Springer-Verlag, 1997.[46] R.F. Egerton, Electron Energy-Loss Spectroscopy in the Electron Microscope,

Plenum Press, New York, 1996.