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A national laboratory of the U.S. Department of Energy 

Office of Energy Efficiency & Renewable Energy 

National Renewable Energy Laboratory 

Innovation for Our Energy Future 

Fundamental Materials Research Subcontract Report 

NREL/SR-520-40568 and Advanced Process September 2006 

Development for Thin-Film 

CIS-Based Photovoltaics 

Final Technical Report 

2 October 2001 — 30 September 2005 

T.J. Anderson, S.S. Li, O.D. Crisalle, 

and V. Craciun 

University of Florida 

Gainesville, Florida 

NREL is operated by Midwest Research Institute Battelle Contract No. DE-AC36-99-GO10337 
Subcontract Report 

NREL/SR-520-40568

September 2006 

Fundamental Materials Research 

and Advanced Process 

Development for Thin-Film 

CIS-Based Photovoltaics 

Final Technical Report 

2 October 2001 — 30 September 2005 

T.J. Anderson, S.S. Li, O.D. Crisalle, 

and V. Craciun 

University of Florida 

Gainesville, Florida

NREL Technical Monitor: B. von Roedern 

Prepared under Subcontract No. ADJ-2-30630-13 

National Renewable Energy Laboratory

1617 Cole Boulevard, Golden, Colorado 80401-3393 

303-275-3000 • www.nrel.gov

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Office of Energy Efficiency and Renewable Energy

by Midwest Research Institute • Battelle 

Contract No. DE-AC36-99-GO10337
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TABLE OF CONTENTS 

1 Development and Characterization of Buffer Layers for Copper-IndiumDiselenide-Based Solar Cells.............................................................................................1

1.1 Introduction..............................................................................................................1 

1.2 Processing ................................................................................................................2 

1.2.1 Chemical-Bath-Deposition Conditions....................................................... 2 

1.2.2 Device Fabrication...................................................................................... 2 

1.3 Characterization of CdS, (Cd,Zn)S, ZnS, and In(OH)xSy Buffer Layers ...............3 

1.3.1 Compositional Analysis.............................................................................. 3 

1.3.2 Optical Properties ....................................................................................... 7 

1.4 Current-Voltage Characteristics of CIGS-Based Solar Cells with CdS,

(Cd,Zn)S, ZnS, or In(OH)xSy Buffer Layers............................................................7 

1.4.1 CIGS-Based Cells with CdS or (Cd,Zn)S Buffer Layers ........................... 7 

1.4.2 CIGS-Based Cells with ZnS Buffer Layers................................................ 9 

1.4.3 CIGS-Based Cells with In(OH)xSy Buffer Layers.................................... 10 

1.5 Conclusions............................................................................................................13 

1.6 References cited.....................................................................................................14 

2 Growth and Characterization of Chemical-Bath-Deposition of ZnxCd1-xS 

Buffer Layers Grown on CIGS Absorbers....................................................................15

2.1 Accomplishments...................................................................................................15 

2.2 Introduction............................................................................................................15 

2.3 Experimental..........................................................................................................16 

2.3.1 Definition of the Relative Zinc Content in the Buffer film and in 

the CBD Solution...................................................................................... 16 

2.3.2 Characterizations of ZnxCd1-xS Thin Films.............................................. 16 

2.3.3 Photo- J-V and QE of CIGS Solar Cells with ZnCdS Buffer Layers ....... 24 

2.4 Conclusions............................................................................................................30 

2.5 References Cited ....................................................................................................31 

3 Effects of Buffer Layer Processing on CIGS Thin Films Characterized by the 

Dual-Beam Optical Modulation Technique...................................................................32

3.1 Introduction............................................................................................................32 

3.2 Simplified Theory for Reflection-Mode DBOM ...................................................33 

iii 
 

3.3 Experimental Details..............................................................................................37 

3.3.1 Lifetime Measurement.............................................................................. 37 

3.3.2 Buffer Layer Processing ........................................................................... 38 

3.4 Results and Discussion ..........................................................................................40 

3.4.1 Measurements of NREL Samples............................................................. 40 

3.4.2 Measurements of SSI and EPV Samples.................................................. 52 

3.5 Conclusions............................................................................................................53 

3.6 References cited.....................................................................................................54 

4 Epitaxial Growth of CuInSe2 at High Temperature.....................................................55

4.1 Introduction............................................................................................................55 

4.2 Experimental Approach .........................................................................................55 

4.3 Composition, Phase and Surface Morphology of CIS Films.................................56 

4.3.1 EPMA Analysis ........................................................................................ 60 

4.3.2 Cross-sectional SEM Images.................................................................... 63 

4.4 XRD Analysis of CIS Films Grown at High Temperature ....................................64 

4.4.1 XRD θ−2θ Scan and ω-rocking Curve Analysis of CIS Films 

Grown at High Temperature................................................................ 64 

4.4.2 Strain Analysis of CIS Films Grown at High Temperature...................... 66 

4.5 TEM Analysis of CIS Films Grown at High Temperature ....................................68 

4.5.1 TEM Analysis of the Cu-rich, Na-free CIS Film (W558)........................ 69 

4.5.2 TEM Analysis of In-rich, Na-free CIS Film (W560) ............................... 72 

4.5.3 TEM analysis of the Cu-rich, Na-dosed CIS film (W584)....................... 76 

4.5.4 TEM Analysis of In-rich, Na-dosed CIS film (W590)............................. 80 

4.6 Conclusions............................................................................................................83 

4.7 References cited.....................................................................................................84 

5 Study of the Defect Structure of CGS by Photoluminescence Spectroscopy .............85

5.1 Introduction............................................................................................................85 

5.2 Theoretical Background.........................................................................................85 

5.3 Experimental Approach .........................................................................................88 

5.4 PL Measurements of Epitaxial CGS Films............................................................89 

5.4.1 PL Spectra................................................................................................. 89 

5.4.2 Thermal Quenching of PL Peaks.............................................................. 96 

5.5 Conclusions..........................................................................................................100 

5.6 References cited...................................................................................................101 

iv 
6 Investigation of Deep-Level Defects in Cu(In,Ga)Se2 Solar Cells by DeepLevel Transient Spectroscopy.......................................................................................103

6.1 Introduction..........................................................................................................103 

6.2 DLTS Principles and Experimental Procedures ..................................................103 

6.3 Results and Discussions.......................................................................................105 

6.3.1 UF Samples from Binary Bilayer Rapid Thermal Process..................... 105 

6.3.2 EPV CIGS Sample.................................................................................. 106 

6.3.3 NREL CIGS Sample from Three Stage Process .................................... 106 

6.4 Conclusions..........................................................................................................107 

6.5 References cited...................................................................................................108 

7 Thermodynamic Modeling of the Isothermal 500°C Section of the Cu-In-Se

System Using a Defect Model........................................................................................129

7.1 Introduction..........................................................................................................129 

7.2 Experimental Phase Equilibria Data in the Cu-In-Se System..............................129 

7.3 Thermodynamic Modeling of Defect-free Cu-In-Se System at 773 K................129 

7.3.1 Gibbs Energy Data of Defect-free α-CIS (CuInSe2) βR-CIS

(Cu1In3Se5), γ-CIS(Cu1In5Se8), and βH-CIS(Cu2In4Se7)......................... 129 

7.3.2 Gibbs Energy of Other Phases................................................................ 130 

7.3.3 Results and Discussion ........................................................................... 131 

7.4 Defect Model of Off-Stoichiometry α-CuInSe2 ..................................................132 

7.4.1 Modeling of Gibbs Energy ..................................................................... 132 

7.4.2 Calculation of the Carrier Concentration in α-CIS (CuInSe2 )............... 137 

7.4.3 Results and Discussion ........................................................................... 139 

7.5 References cited...................................................................................................139 

8 Formation of α-CuInSe2 by Rapid Thermal Processing of a Stacked Binary

Compound Bilayer.........................................................................................................156

8.1 Introduction..........................................................................................................156 

8.2 Background..........................................................................................................156 

8.3 Experimental apparatus........................................................................................157 

8.4 Film Characterization...........................................................................................158 

8.4.1 Crystal Structure Determination by X-ray Diffraction (XRD)............... 158 

8.4.2 Topographical Information - Scanning Electron Microscopy

(SEM) ..................................................................................................... 158 

8.4.3 Local Composition Analysis - Energy Dispersive Spectroscopy

(EDS) ...................................................................................................... 159 

v 
8.4.4 Composition Depth Profile - Auger Electron Spectroscope (AES)........ 159 

8.4.5 Local Composition Analysis - Electron Probe X-ray Microanalysis 

(EPMA)................................................................................................... 159 

8.4.6 Overall Composition Analysis - Inductively Coupled Plasma (ICP) ..... 159 

8.4.7 I-V Measurements .................................................................................. 159 

8.5 Experimental Results and Discussions ................................................................160 

8.5.1 Susceptor Design .................................................................................... 160 

8.5.2 Phase Transition Study on CuSe/InSe/Mo/Glass Precursors ................. 163 

8.5.3 RTP Experiments on CuSe/InSe/Mo/Glass Precursors in PVD 

system ..................................................................................................... 165 

8.5.4 RTP Experiments on InSe/CuSe/Mo/Glass Precursors .......................... 167 

8.6 Conclusions..........................................................................................................167 

8.7 References cited...................................................................................................167 

9 Investigation of Pulsed Non-Melt Laser Annealing on the Film Properties 

and Performance of Cu(In,Ga)Se2 Solar Cells. ..........................................................198

9.1 Accomplishments.................................................................................................198 

9.2 References cited...................................................................................................205 

10 Device Modeling and Simulation of CIGS Solar Cells ...............................................206

10.1 Accomplishments.................................................................................................206 

10.2 Introduction..........................................................................................................206 

10.3 Cell Structure and Material Parameters ...............................................................207 

10.4 Effect of Graded Band-gaps on Cell Performance ..............................................208 

10.5 Comparison with Experimental Data...................................................................212 

10.6 Conclusions..........................................................................................................216 

10.7 References Cited ..................................................................................................216 

vi 
1 Development and Characterization of Buffer Layers for 

Copper-Indium-Diselenide-Based Solar Cells 

1.1 Introduction 

In a typical CIGS-based solar cell a very thin CdS buffer layer (< 500Å) is usually 

deposited between the ZnO window layer and the CIGS absorber layer in order to achieve high 

conversion efficiency. Among various buffer layer materials such as CdS, (Cd,Zn)S, ZnS, 

Zn(O,S,OH)x, ZnO, ZnSe, Inx(OH,S)y, In2S3, In(OH)3, SnO2, Sn(S,O)2, ZnSe, or ZrO2 deposited 

by the CBD, ALE, MOCVD, or other deposition processes, the best performance CIGS solar cell 

with a total-area conversion efficiency of 18.8% [1-1] and other high-efficiency (> 17%) CIGS 

solar cells [1-2][1-3] were obtained by using the CBD CdS buffer layers. However, the role of 

CBD CdS buffer layer on the CIGS cell performance is not well understood. In view of the 

ability of depositing large area and uniform films, the non-vacuum low-temperature CBD 

process is very advantageous for manufacturing low-cost photovoltaic devices. Although a dry 

vacuum process has advantages for in-line manufacturing, the CBD process provides a simple, 

flexible, and cost-effective means for the deposition of buffer layers on the CIS-based solar cells. 

It has been reported that Cd-partial electrolyte treatment [1-4] modifies the CIGS surface 

favorably as evidence by the improvement in performance of CIGS solar cells [1-5]. 

Unfortunately, the basic mechanisms responsible for the cell performance are still not clear. 

Using the wider band-gap materials to replace the CdS (Eg≈2.4eV) buffer layer could 

improve the quantum efficiency of the CIGS cell at shorter wavelengths, resulting in an increase 

of the short-circuit current. The (Cd,Zn)S buffer layer has a band-gap energy greater than 2.4eV, 

and can provide a better lattice match to the CIGS absorber layer. On the other hand, the toxicity 

of cadmium (Cd) and the desirability of using a wider band-gap material to achieve a higher 

spectral response in the blue region have motivated the search for other alternative buffer layer 

materials. Using ZnS (Eg≈3.6eV) and In(OH)xSy (Eg≈2.54eV [1-6]) buffer layers for the CIGS 

solar cells have achieved high active-area conversion efficiencies of 18.1% [1-7] and 15.7% [1­

8], respectively. Thus, both ZnS and In(OH)xSy thin films are promising candidates for the Cdfree buffer layers among the reported alternative materials 

The buffer layer materials of CdS, (Cd,Zn)S, ZnS, or In(OH)xSy was deposited on the 

soda-lime glass substrates, Cu(In,Ga)Se2 (CIGS), or Cu(In,Ga)(Se,S)2 (CIGSS) thin films by the 

chemical-bath-deposition (CBD) process. The impurities in the deposited films and their atomic 

concentration were characterized by the x-ray photoelectron spectroscopy (XPS) and auger 

electron spectroscopy (AES) analyses. In addition, the composition of depth profile of the 

samples was characterized by using the AES. The band-gap energy of the deposited In(OH)xSy 

thin films was determined from the optical absorption data. Both the CIGS and CIGSS samples 

deposited with the CdS, (Cd,Zn)S, ZnS, or In(OH)xSy buffer layers by the CBD process were 

fabricated into solar cells. The current-voltage (I-V) characteristics of these cells with the 

alternative buffer layers were measured, and the results were compared to the cells deposited 

with CBD CdS buffer layers. The results show a comparable performance among these cells, 

and further optimization of the deposition conditions should improve the performance of the 

cells deposited with the alternative buffer layers. 

1
1.2 Processing 

1.2.1 Chemical-Bath-Deposition Conditions 

In the deposition of thin solid films in a chemical bath by the CBD process the nucleation 

centers are regularly formed by the absorption of metal hydroxo species on the surface of the 

substrate. The initial layer of the thin film is formed through the replacement of hydroxo group 

by the sulphide ions, and subsequently the solid film is grown by the condensation of metal and 

sulphide ions onto the top of the initial layer. 

The aqueous solution consisting of 2.4×10-4M CdCl2, 2.38×10-3M NH2CSNH2, 7.43×10­

4

M NH4Cl, and 6.61×10-4M NH4OH at a bath temperature in the range of 80 to 85°C was applied 

for the deposition of CdS films. The CBD In(OH)xSy films were deposited using a freshly 

prepared aqueous solution of 0.005M indium chloride and 0.15M thioacetamide at 70oC with 

deposition times varied between 15 and 25 minutes and with a pH value of 1.8. In order to avoid 

nucleation centers in the bath for the deposition of In(OH)xSy, the insoluble particles present in 

the stock solution of thioacetamide were eliminated by filtration. The complexing agent, acetic 

acid, was added into the bath during the deposition of In(OH)xSy as a new deposition condition 

to improve the film quality. The deposition of (Cd,Zn)S buffer layers was carried out with 

1.2×10-3M CdCl2, 6.27×10-4M ZnCl2, 1.2×10-2M thiourea, 5.27×10-4M NH3, and 1.39×10-3M 

NH4Cl, and the bath temperature was maintained at around 85°C. For the CBD ZnS films the 

deposition bath was prepared with an aqueous solution of 2.5×10-2M ZnSO4, 3.5×10-2M thiourea, 

1M NH3, and 3M hydrazine at a bath temperature varied between 70 and 80°C. 

Two competitive processes, namely the heterogeneous process of the solid film deposited 

on the substrate and the homogeneous process of precipitation in the reaction bath, took place 

simultaneously in the chemical bath during the CBD process. In order to remove the possible 

precipitation, which may be attached on the surface of the films, the substrates were well rinsed 

with DI water after the deposition and blown dry with nitrogen stream. 

1.2.2 Device Fabrication 

The CdS, (Cd,Zn)S, ZnS, or In(OH)xSy buffer layers were deposited on the CIGS and 

CIGSS samples provided by the Institute of Energy Conversion (IEC) of the University of 

Delaware, ISET, and Siemens Solar Industries. The high/low resistivity ZnO bilayer films were 

then deposited by either the MOCVD or sputtering technique on these samples. The Ni/Al grids 

as contact pads were finally deposited to obtain the finished cells. 

2
1.3 Characterization of CdS, (Cd,Zn)S, ZnS, and In(OH)xSy

Buffer Layers 

1.3.1 Compositional Analysis 

 The near surface composition of the CdS, (Cd,Zn)S, ZnS, and In(OH)xSy films deposited 

on the soda-lime glass substrates and CIGS samples were characterized by the x-ray 

photoelectron spectroscopy (XPS), and the results are shown in Figure 1-1. The XPS analysis 

was conducted by using a Physical Electronics 5100 ESCA equipped with a Mg Kα X-ray 

source and a hemispherical analyzer. Depth profiling was performed with an argon ion source 

operating at the beam energy of 4 KeV. The charging effects were taken into account by 

referring the measured spectra to the binding energy peak of C 1s at 285eV. Carbon was initially 

present for all films deposited by the CBD process described in the previous section but became 

undetectable during sputter depth profiling, indicating that little carbon was incorporated in the 

films. Besides carbon, elements of metal (Cd, Zn, or In), sulfur, and oxygen were also identified 

in the deposited films. 

Auger 

(b)(Cd,Zn)S OAuger CdCd O1s

Cd3d3/2 

(405.5eV) 

C S2p (532eV) 

(a) CdS Cd Cd3d5/2 

Zn2p3/2 (285eV) S (162eV) Cd 

Auger 

Zn2p3/2 

Cd

OAuger Cd Cd O1s

Cd3d3/2 Cd3d5/2 Zn 

(1022eV) (532.5eV) (405.5eV) C S2p 

O1s 

(285eV) S (162eV) Cd 

(c) ZnS OAuger (531.5eV) ZnAuger C Zn

(285eV) Auger S2p InAuger In3d3/2 In3d5/2 (162eV) Zn Zn 

(d) In(OH)xSy 

OAuger In In O1s (444.5eV)

C S2p (531.5eV) 

(285eV) S (161.5eV) In 

1200 1100 1000 900 800 700 600 500 400 300 200 100 0 

Binding Energy (eV) 

Figure 1-1. XPS spectra of CdS, (Cd,Zn)S, ZnS, or In(OH)xSy thin films deposited 

on CIGS thin films by the CBD process. 

 Before and after sputtering the CBD CdS films the binding energy peaks of Cd 3d5/2 and 

S 2p3/2 at about 405.3eV and 161.5eV were obtained, respectively. The result shows that the 

sulfur peak is in good agreement with the reported data (S-Cd: 161.7eV [1-9]). The binding 

energies of Cd-S (405.3eV [1-9]), Cd-O (405.2eV [1-9]), and Cd-OH (405.0eV [1-9]) are too 

close to identify the exact compounds presented in the film from the measured Cd peak at 

405.3eV. Two peaks at 529.48eV and 531.09eV obtained from the deconvolution of O 1s 

binding energy curve demonstrate that the film consists of chemical bonds Cd-O (529.2eV [1-9]) 

3 
and Cd-OH (530.9-532eV [1-9]). The estimated atomic concentration ratios of [Cd]/[S] and 

[S]/[O] are approximately 1.6 and 5.4, respectively. For the (Cd,Zn)S film the atomic 

concentration was estimated from the peak area of each element, Cd, Zn, S, and O, in the film. 

The XPS result indicated that a large amount of oxygen (around 40%) and a small amount of 

zinc (around 2%) were incorporated in the (Cd,Zn)S film. Because the precipitation of CdS is 

easier and faster than that of ZnS during the CBD process, the deposited (Cd,Zn)S film only had 

a very small quantity of zinc even though the concentration of zinc source was an order of 

magnitude higher than that of cadmium source in the deposition bath. 

For the ZnS film, the deconvolution of Zn 2p3/2 binding-energy spectrum from 1018eV to 

1026eV revealed only one peak at 1022.0eV. This is consistent with the energy peaks of 

possible compounds ZnS and/or ZnO at 1022.0eV and 1021.8eV-1022.5eV [1-9], respectively. 

No evidence of ZnSO4 was found from the Zn 2p3/2 peak, where the binding energy peak of 

ZnSO4 is 1023.1eV. The oxygen 1s photoelectron binding energy spectra were fitted with 

Gaussian-Lorentzian curves both before and after sputtering the film. The XPS spectra can be 

represented by two peaks at about 530.6-530.1eV and 531.6eV, which correspond to the 

chemical bonds of Zn-O (530.4eV [1-9]) and metal-hydroxide compound (530.9-532eV [1-9]), 

respectively. Thus the results suggest that the CBD ZnS films might be a mixture of ZnS, ZnO, 

and Zn(OH)2. The estimated atomic concentration ratio of Zn, S, and O was about 5:3:2 after 

sputtering the film for 4 minutes. 

 For the In(OH)xSy films deposited by CBD with acetic acid in the bath, a binding energy 

peak of S 2p3/2 was found at about 161.24eV (see Figure 1-2), which can be assigned to the 

sulfide compound. In Figure 1-2, a very weak binding energy peak at around 168.96eV suggests 

that the indium sulfate (169-171eV [1-9]) was also present. After sputtering the film for 4 

minutes no peak was found between 167ev and 173eV. We conclude that the sulfate compound 

only exists at the surface of the film. From Figure 1-3 the montage of binding energy 

distribution curves between 526eV and 544eV, it is evident that the shape of these curves 

changes for the oxygen 1s peak during the depth profiling. As shown in the inset of Figure 1­

3(a), before sputtering the XPS spectra can be represented by two peaks at 531.4eV and 

529.5eV, suggesting that the oxygen is present in the form of In-OH (531.8eV [1-9]) and In-O 

(529.8eV [1-9]). As exhibited in the inset of Figure 1-3(b), likewise two fitted peaks at 531.5eV 

and 530.0eV are present in the curve after sputtering the film. This suggests that the film has 

more indium hydroxide than indium oxide in the near surface region. The estimated atomic 

concentration ratios of In, O, and S were found to be around 49.9%: 26.5%: 23.6% after 

sputtering for 4 minutes. 

4
Before sputtering 

Binding Energy (eV) 

171 169 167 165 163 161 159 

After sputtering 

171 169 167 165 163 161 159 

Before sputtering

Binding Energy (eV)

171 169 167 165 163 161 159

After sputtering

171 169 167 165 163 161 159

Figure 1-2. XPS spectra of In(OH)xSy films. 

(a) (b) 

531.4eV 

529.5eV 531.5eV 

530.0eV 

(a) (b)

531.4eV

529.5eV 531.5eV

530.0eV

542 538 534 530 526 542 538 534 530 526

Binding Energy (eV)

Figure 1-3. Montage of photoelectron binding energy spectra for a depth profile 

of an In(OH)xSy film (View with surface curve in the background). The oxygen 1s 

photolectron binding energy spectra are (a) as-received surface, and (b) after 

sputtering for 4 minutes. 

To determine whether the carbon was only present at the surface or incorporated 

throughout the film, Auger Electron Spectroscopy (AES) sputter depth profiles were performed. 

Figure 1-4 shows the differential Auger spectra taken for the initial In(OH)xSy film and 

compared to those taken after 3 and 5 minutes of sputtering. The results indicate that the carbon 

is only at the surface and likely derived from the ambient. The spectra in Figure 1-4 also show 

that after 5 minutes sputtering time, some regions of the glass were exposed as evidenced by the 

appearance of Ca and Si signals. An AES depth profile for the CIGS/In(OH)xSy sample is shown 

in Figure 1-5. The results indicate that the buffer layer barely covered the CIGS layer. Although 

intermixing between the In(OH)xSy and the CIGS layers is suggested from the depth profile, this 

result is probably attributed to the roughness of CIGS surface. Interdiffusion between the 

In(OH)xSy buffer layer and the CIGS absorber layer is not expected to occur during the low 

temperature CBD process. However, an observation of intermixing at the interface between CdS 

5
and CIGS involving the elements of S, Se, and In was reported recently [1-10]. Furthermore 

after the completed cell processing a strong intermixing between absorber (CIGS), buffer (CdS), 

and TCO (ZnO) material in the interface region was found [1-11]. 

30 

initially 

25 C O 

(270eV) In 

20 S (407eV) (513eV) 

(151eV) 

after sputtering time of 3 minutes 

15 

10 

after sputtering time of 5 minutes 

5 

0 

Ca

Si (289eV) (74eV) 

S In 

(150eV) (402eV) 

O 

(504eV) 

0 100 200 300 400 500 600 700 800 900 1000 1100 

Kinetic Energy (eV) 

Figure 1-4. Differential Auger spectra of the In(OH)xSy thin film deposited on the 

glass substrate. 

60 

Se 

50 

40 

S 

30 Cu 

20 O

In 

10 Ga 

0 

0 5 10 15 20 25 30 

Sputtering time (minutes) 

Figure 1-5. The Auger depth profile of the In(OH)xSy thin film deposited on the 

CIGS cell by CBD process. 

6 
1.3.2 Optical Properties 

The spectral dependence of the transmittance for the In(OH)xSy films grown on the sodalime glass substrates was measured using a spectrophotometer, and the results are shown in the 

inset of Figure 1-6. The In(OH)xSy films with shorter deposition time have a higher 

transmittance at shorter wavelengths, which will result in an increase of short-circuit current in 

the CIGS cells. The buffer layers grown with longer deposition time, however, could provide a 

better shield effect against the sputtering damage during the ZnO deposition step and thereby 

reduce the interface recombination. Thus, the buffer-layer thickness needs to be optimized for 

the solar-cell performance. The band-gap energy of In(OH)xSy was determined from the optical 

absorption data. To obtain the direct transition, (αhν)

2

 was plotted against hν for the In(OH)xSy 

films as illustrated in Figure 1-6. A band-gap energy of 2.54eV was obtained from the intercept 

of this plot. The band-gap energy of In(OH)xSy, which depends on the films stoichiometry, lies 

between 2eV and 3.7eV, the band-gap energies of In2S3 and In2O3, respectively [1-12]. 

3 100 

90 thickness=43nm 

2.5 70 

80 

57nm 

60 76nm 

2 50 

40 

30 

1.5 20 

10 

0 

1 300 400 500 600 

wavelength (nm) 

700 800 

0.5 

0 

2.3 2.4 2.5 2.6 2.7 2.8 

hν (eV) 

Figure 1-6. (αhν)

2

 versus hν plot for the In(OH)xSy thin film deposited on the 

glass substrate by the CBD process. 

1.4 Current-Voltage Characteristics of CIGS-Based Solar Cells 

with CdS, (Cd,Zn)S, ZnS, or In(OH)xSy Buffer Layers 

1.4.1 CIGS-Based Cells with CdS or (Cd,Zn)S Buffer Layers 

The J-V characteristics of CIGSS solar cells with the CdS or (Cd,Zn)S buffer layers 

deposited by our baseline CBD process are shown in Figure 1-7, where the CIGSS absorbers are 

provided by the Siemens Solar Industries. Evidently the CIGSS solar cell with the CdS buffer 

7
 Eff. Voc (V) Jsc (mA/cm2

) F.F. 

(a) CdS 11.1% 0.53 31.4 66.3% 

(b) (Cd,Zn)S 9.5% 0.51 29.5 61.7% 

(c) ZnS 9.8% 0.52 27.7 66.4% 

(d) In(OH)xSy 8.6% 0.525 27.3 59.7% 

(a) (c) 

(d) 

(b) 

layer has the best conversion efficiency. On the contrary the performance of the other cells 

deposited with the (Cd,Zn)S, ZnS, or In(OH)xSy buffer layers yields a lower conversion 

efficiency than that of the CIGSS/CdS cell. From the results of this study further optimization of 

the alternative buffer layers is needed, which includes the optimization of the deposition 

conditions to improve the film quality and the investigation of the effects of post-deposition 

annealing and light soaking on the devices. 

0 

5 

10 

15 

20 

25 

30 

35 

Current Density (mA/cm2 ) 

0

5

10

15

20

25

30

35

Current Density (mA/cm2)

0 0.1 0.2 0.3 0.4 0.5 0.1 0.2 0.3 0.4 0.5

Voltage (V)Voltage (V)

Figure 1-7. The J-V characteristics of the CIGSS cells with CdS, (Cd,Zn)S, ZnS, 

or In(OH)xSy buffer layers. 

In addition, the CdS buffer layers were also deposited by using the UF baseline CBD 

process [1-13] with a deposition time of 30 or 40 minutes on the same sets of CIGS samples 

provided by the IEC. The CIGS/CdS cell with a deposition time of 30 minutes has a better 

performance with VOC =0.51V, JSC =30.5mA/cm2, F.F. =63.8%, and conversion efficiency η

=9.99%, and the results are listed in Table 1-1. Apparently, the effect of annealing considerably 

improves the performance of these cells with either the thinner or thicker buffer layers. Also the 

performance of the CIGS cells can be further enhanced by optimizing the thickness of the CdS 

buffer layers via the variation of the deposition time. 

8
Table 1-1. The performance of CIGS Cells with CBD CdS buffer layers 

 1.1.1 Deposition 

time 

As deposited Annealing 5 min. 

Conversion Efficiency η

40minutes 8.87% 9.91% 

30minutes 8.41% 9.99% 

Fill Factor (F.F) 

40minutes 61.7% 61.9% 

30minutes 60.8% 63.8% 

Open-circuit Voltage VOC 

40minutes 0.4745V 0.5017V 

30minutes 0.4697V 0.5136V 

Short-circuit Current Density 

JSC 

40minutes 30.31 

mA/cm2 

31.88 mA/cm2 

30minutes 29.47 

mA/cm2 

30.51 mA/cm2 

1.4.2 CIGS-Based Cells with ZnS Buffer Layers 

Figure 1-8 illustrates the effect of ZnS buffer-layer thickness on the performance of 

Siemens CIGSS solar cells. By increasing the ZnS buffer-layer thickness the absorption of 

incident light in the ZnS layer increases, and thus the short- circuit current of the cell decreases. 

Due to the high resistivity of ZnS buffer layer a thick buffer layer can cause a high series 

resistance in the solar cell and results in a poor fill factor as shown in Figure 1-8. With a higher 

band gap of ZnS than that of CdS, the CIGSS/ZnS cell should have a higher short-circuit current 

than the CIGSS/CdS cell because of the higher quantum efficiency at the shorter wavelength 

region. However, the highest short-circuit current obtained in this study for the CIGSS/ZnS cell 

was only 27.74 mA/cm2 (see Figure 1-7), whereas for the CIGSS/CdS cell a short-circuit current 

of 31.4 mA/cm2

 was obtained. Thus, the further optimization of ZnS layer thickness and 

deposition condition is needed in order to achieve the potential of ZnS buffer layers for the CISbased cells. The performance of the CIGS solar cells fabricated from the same set of ISET CIGS 

samples with our ZnS or ISET CdS buffer layers is shown in Figure 1-9. The CIGS/ZnS cell has 

a higher short-circuit current as expected but the cell suffers a comparably low open-circuit 

voltage. 

9
0 

5 

10 

15 

20 

25 

30 

(c) 

(b) 

(a) 

Current Density (mA/cm2 ) 0

5

10

15

20

25

30

(c)

(b)

(a)

Current Density (mA/cm2)

0 0.1 0.2 0.3 0.4 0.5 0.6 0 0.1 0.2 0.3 0.4 0.5 0.6

Voltage (V)Voltage (V)

Figure 1-8. The J-V characteristics of CIGSS solar cells using different thickness 

of CBD ZnS buffer layers deposited for (a) 10min, (b) 20min, and (c) 25min at 

70°C. 

0 

5 

10 

15 

20 

25 

30 

35 

40 (b) CBD ZnS Buffer layer 

(a) ISET CdS Buffer layer 

Eff. Voc (V) Jsc (mA/cm2) F.F. 

(a) CdS 9.9% 0.435 33.6 68.1% 

(b) ZnS 9.3% 0.396 36.6 64.1% 

Current Density (mA/cm2 ) 

0

5

10

15

20

25

30

35

40 (b) CBD ZnS Buffer layer

(a) ISET CdS Buffer layer

Eff. Voc (V) Jsc (mA/cm2) F.F.

(a) CdS 9.9% 0.435 33.6 68.1%

(b) ZnS 9.3% 0.396 36.6 64.1%

Current Density (mA/cm2)

0 0.05 0.1 0.15 0.2 0.25 0.3 0.35 0.4 0.45 0 0.05 0.1 0.15 0.2 0.25 0.3 0.35 0.4 0.45

Voltage (V)Voltage (V)

Figure 1-9. The J-V characteristics of CIGS cells with CdS buffer layer deposited 

at ISET or the ZnS buffer layer deposited at UF. 

1.4.3 CIGS-Based Cells with In(OH)xSy Buffer Layers 

In order to obtain an optimum CBD process procedure for the deposition of In(OH)xSy 

buffer layers, three different sequences of adding chemicals InCl3 and CH3CSNH2, which might 

result in the modification of CIGS surface, were applied to the CIGS samples. It was found that 

the CIGS sample dipped in the DI water for a few seconds prior to the start of the deposition, 

which might have a precleaning effect on the CIGS surface, showed the best device performance 

from this experiment. The results are shown in the Table 1-2. 

10
 1.1.1 Cell- 1 Cell- 2 Cell- 3 

Conversion Efficiency η 7.39 % 5.73 % 5.23 % 

Fill Factor (F.F) 44.6 % 38.2 % 38.3 % 

Open-circuit Voltage VOC 0.57 V 0.52 V 0.51 V 

Short-circuit Current Density JSC 29.1 mA/cm2 28.9 mA/cm2 27.0 mA/cm2 

Table 1-2. The Performance of the CIGS solar cells with CBD In(OH)xSy buffer 

layers using different procedures of adding chemicals 

Note: For the CBD process: cell- 1 was first immersed in DI water, and then both InCl3 and 

CH3CSNH2 were added into the bath simultaneously; cell- 2 was first immersed in the solution of 

CH3CSNH2, and then InCl3 was added into the bath; cell- 3 was first immersed in the solution of 

InCl3, and then CH3CSNH2 was added into the bath. 

Each I-V measurement for the CIGS cells comprised a sweep in the ISC to VOC direction 

followed immediately by a sweep from VOC to ISC direction. The dark- and photo-I-V curves of 

the CIGS/In(OH)xSy heterojunction cell before and after annealing at 200 °

C in the air for 20 

minutes are shown in the Figure 1-10. As illustrated a hysteresis loop was observed in the photoI-V curves of the CIGS/In(OH)xSy cells before and after annealing. On the contrary, no evidence 

of a hysteresis loop in the photo-I-V curves of the CIGS/CdS cells was found in this study. 

30 

10 

20 before annealing 

after annealing 

0 

-10 

-20 

-30 

-40 

-0.5 -0.4 -0.3 -0.2 -0.1 0 0.1 0.2 

Voltage (V) 

0.3 0.4 0.5 0.6 0.7 

Figure 1-10. The J-V characteristics of the CIGS cell deposited with CBD 

In(OH)xSy buffer layer before and after annealing. 

11
Initially the device without annealing showed the inflected I-V curve. As shown in 

Figure 1-10 and Figure 1-11, the fill factor (F.F.), JSC, VOC, and hence the conversion efficiency 

were dramatically improved after annealing time of 20 minutes. The cells, however, showed 

degradation when the total annealing time was increased to 40 minutes. In order to improve the 

cell performance, it is necessary to optimize the annealing conditions. 

Figure 1-11. The performance of the CIGS solar cells deposited with CBD 

In(OH)xSy buffer layers as a function of annealing time. 

Comparing the performance of CIGS/CdS cells and CIGS/In(OH)xSy cells, we have 

found that the latter has a much higher open-circuit voltage, a comparable short-circuit current 

density, and a very poor fill factor (VOC =0.57V, JSC =29.1mA/cm2

, F.F. =44.6%, and a low 

conversion efficiency η =7.39%) as shown in Figure 1-12, while the CIGS/CdS cell yielded a 

conversion efficiency of 9.99%, VOC=0.51V, JSC=30.5mA/cm2

, and F.F.=63.8%. The possible 

reason why the CIGS/In(OH)xSy cell has a higher VOC is that the net acceptor concentration in 

the CIGS layer is increased and hence the depletion width is reduced. The CBD alternative 

buffer layer process might influence the charged states at the surface or at the grain boundaries of 

the CIGS layer and hence modify the electrical properties of the absorber layer as well [1-8]. 

However, the CIGS/In(OH)xSy cell has a lower JSC due to the lower quantum efficiency than the 

CIGS/CdS cell in the long wavelength region [1-8]. 

12
-40 

-30 

-20 

-10 

0 

10 

20 

30 

CdS buffer layer 

In 

x

(OH,S)y buffer layer 

-0.1 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 

Voltage (V) 

Figure 1-12. J-V characteristics of the CIGS cells with CBD- CdS or CBDIn(OH)xSy buffer layer (CIGS/CdS solar cell: Eff.=9.99%, VOC=0.51V, 

JSC=30.5mA/cm2, F.F.=63.8%; CIGS/In(OH)xSy solar cell: Eff.=7.39%, 

VOC=0.57V, JSC=29.1mA/cm2, F.F.=44.6%). 

By adding the complexing agent of acetic acid into the chemical bath, a more adherent, 

reproducible, homogeneous, and higher quality In(OH)xSy film [1-14] has been deposited on the 

glass substrates and Siemens CIGSS samples. The performance of one CIGSS cell deposited 

with In(OH)xSy buffer layer using acetic acid and the performance of the CIGSS cells deposited 

with other buffer layers are shown in Figure 1-7. The result reveals that we can take advantage 

of the In(OH)xSy buffer layer deposited by the new deposition method to improve the cell 

performance. 

1.5 Conclusions 

A study of the electrical and optical properties of the CBD buffer layers on the soda-lime 

glass substrates and on the absorber layers was conducted. The XPS results indicated that CdS, 

(Cd,Zn)S, ZnS, and In(OH)xSy films deposited by the CBD method were not completely pure. A 

small amount of carbon was found to incorporate in the first few atomic layers of the films. Also 

oxygen in the form of metal oxide and/or metal hydroxide was incorporated into the films during 

the deposition process. A better quality In(OH)xSy buffer layer has been achieved by adding the 

acetic acid into the reaction bath. Annealing has been found to be a key factor for improving the 

conversion efficiency of CIS-based cells deposited with alternative buffer layers. Among the 

CIS-based solar cells with CdS, (Cd,Zn)S, ZnS, or In(OH)xSy buffer layers studied in this work, 

the CIGSS/CdS cell gave the best performance. From the J-V characteristics obtained for the 

CIS-based solar cells with the alternative buffer layers, the deposition conditions of CBD process 

are merited to further optimize in order to realize the potential advantages of these buffer layers. 

Future study of the CIS-based solar cells with alternative buffer layers will focus on optimizing 

annealing conditions and on investigating the light soaking effects. 

13
1.6 References cited 

[1-1] M. Contreras, B. Egass, K. Ramanathan, J. Hiltner, A. Swartzland, F. Hasoon, and R. 

Noufi, Progress in Photovoltaics, Vol. 7, pp. 311, 1999. 

[1-2] T. Negami, Y. Hashimoto, and S. Nishiwaki, Solar Energy Materials and Solar Cells, Vol. 

67, pp. 331-335 ,2001. 

[1-3] E. Niemi, J. Hedstrom, T. Martinsson, K. Granath, L. Stolt, J. Skarp, D. Hariskos, M. 

Ruckh, H.W. Schock, Conference Record of the 25th IEEE Photovoltaic Specialists 

Conference, pp. 801-804, 1996. 

[1-4] J. Kessler, K.O. Velthaus, M. Ruckh, R. Laichinger, H.W. Schock, D. Lincot, R. Ortega 

and J. Vedel, 6th Photovoltaic Science and Engineering Conference (PVSEC-6) 

Proceedings, New Delhi, India, pp. 1005-1010, 1992. 

[1-5] K. Ramanathan, R.N. Bhattacharya, J. Granata, J. Webb, D. Niles, M.A. Contreras, H. 

Wiesner, F.S. Hasoon, and R. Noufi, Conference record of the 26th IEEE Photovoltaic 

Specialists Conference, p. 319, 1997. 

[1-6] C.H. Huang, Sheng S. Li, W.N. Shafarman, C.-H. Chang, J.W. Johson, L. Reith, S. Kim, 

B.J. Stanbery, and T.J. Anderson, Technical Digest of the 11th International Photovoltaic 

Science and Engineering Conference, p. 855, 1999. 

[1-7] T. Nakada, and M. Mizutani, Japanese Journal of Applied Physics, Part 2, Vol. 41, Issue 

2B, pp. L165-L167, 2002. 

[1-8] D. Hariskos, M. Ruckh, U. Rühle, T. Walter, H.W. Schock, J. Hedström, and L. Stolt, 

Solar Energy Materials and Solar Cells, Vol. 41/42, pp. 345-353, 1996. 

[1-9] J.F. Moulder, W.F. Stickle, P.E. Sobol, and K.D. Bomben, Handbook of X-ray 

Photoelectron Spectroscopy, pp. 45, 61, and 213-242, 1995. 

[1-10] C. Heske, D. Eich, R. Fink, E. Umbach, T. van Buuren, C. Bostedt, L.J. Terminello, S. 

Kakar, M.M. Grush, T.A. Callcott, F.J. Himpsel, D.L. Ederer, R.C.C. Perera, W. Riedl, 

and F. Karg, Applied Physics Letters, Vol. 74, pp. 1451-1453, 1999. 

[1-11] I. Luck, U. Störkel, W. Bohne, A. Ennaoui, M. Schmidt, H.W. Schock, and D. Bräunig, 

Thin Solid Films, Vol. 387, pp. 100-103, 2001. 

[1-12] R. Bayon, C. Guillen, M.A. Martinez, M.T. Gutierrez, and J. Herrero, J. Electrochem. 

Soc., Vol. 145, pp. 2775-2779, 1998. 

[1-13] C. -H. Chang, A. A. Morrone, B. J. Stanbery, C. McCreary, M. Huang, C. H. Huang, S. S. 

Li, and T. J. Anderson, AIP Conference Proceedings (NCPV Program Review Meeting), 

No. 462, pp. 114-119, 1998. 

[1-14] R. Bayon, C. Maffiotte, and J. Herrero, Thin Solid Films, Vol. 353, pp. 100-107, 1999. 

14
2 Growth and Characterization of Chemical-BathDeposition of ZnxCd1-xS Buffer Layers Grown on CIGS 

Absorbers 

2.1 Accomplishments 

This report presents a characterization of thin films of ZnxCd1-xS (denoted as ZnCdS in 

the suite), prepared by chemical bath deposition (CBD) for deployment as a buffer layer in CIGS 

thin film solar cells. We determined that ZnCdS materials with specific Zn contents have 

attractive physical properties for deployment as buffer films in CIGS solar cells. In particular, 

we deposited buffer layers of various relative stoichiometric compositions of Zn, denoted by the 

subindex x in the chemical formula ZnxCd1-xS, and investigated their structural, surface, optical, 

and electrical properties by XRD, SEM, spectrophotometry, and resistivity studies. The degree 

of crystallinity in the film decreased with increasing relative zinc content x, leading to an 

amorphous structure at the concentration level x = 0.5. In addition, the resistivity of the ZnCdS 

films also increased with increasing relative zinc content x. Of particular relevance is also the 

fact that the energy band gap of ZnCdS is higher than that of CdS, and that the bandgap increases 

with increasing zinc content x. 

We manufactured CIGS solar cells using ZnxCd1-xS buffer layers with Zn content values 

x = 0 (CdS), 0.1, 0.2, 0.3, 0.4, and 0.5, and carried out a comprehensive device characterization 

study in terms of current-density/voltage and quantum efficiency. The results showed that a 

ZnCdS/CIGS cell using buffer layer with a Zn composition x = 0.2 realized a conversion 

efficiency that was greater by a factor of 1.30 than the efficiency of a Zn-free reference cell of 

the form CdS/CIGS. The CIGS absorber films were provided by EPV. 

2.2 Introduction 

In view of environmental safety, it is highly desirable to replace CdS (Eg ≈ 2.4 eV) by 

cadmium (Cd) free and toxic-free alternative buffer layers. In addition, wider band-gap buffer 

materials are preferred for CIGS cells, since they allow transmission of shorter wavelength 

photons into CIGS absorber and hence higher short-circuit current density (Jsc) could be 

expected. For this purpose, zinc (Zn)-based compound materials such as ZnO (Eg ≈ 3.3 eV), 

Znx(O,S)y (Eg ≈ between 3.3 and 3.8 eV), Zn(O,S,OH)x, and ZnS (Eg ≈ 3.8 eV) deposited by 

CBD process are proposed as alternative buffer layers for CIGS thin film solar cells. 

The addition of Zn to the most widely used CdS buffer layer material decreases the lattice 

constant with a lattice-match to CIGS and CGS absorbers and produces a more favorable 

conduction band alignment. Increasing Zn concentration in the ZnxCd1-xS (CdS: Eg ≈ 2.4 eV, 

ZnS: Eg ≈ 3.8 eV) films increases band-gap energy, which results in an increase in short 

wavelength photo-response. Therefore, ZnxCd1-xS could lead to the increase in quantum 

efficiency (QE) at shorter wavelengths and a favorable conduction band offset at the interface 

ZnxCd1-xS buffer and the CIGS or CGS absorber layer. Furthermore, the influence of Zncomposition, x, in the ZnxCd1-xS films on the properties of interface layers has not yet fully 

15
understood. Therefore, studies of ZnxCd1-xS buffer layers to replace the CdS buffer layer for the 

CIGS and CGS solar cells are desirable. 

This report presents the development, growth, and characterization of ZnxCd1-xS (denoted 

as ZnCdS) films prepared by chemical bath deposition (CBD) for use as buffer layers in CIGS 

thin film solar cells and the characteristics of I-V and QE of the CIGS solar cells with various Zn 

compositions in ZnCdS buffer layers. The results are compared to the cells deposited with CBD 

CdS buffer layer. 

2.3 Experimental 

2.3.1 Definition of the Relative Zinc Content in the Buffer film and in the CBD Solution 

The composition of Zn relative to that of Cd in the chemical-bath-deposition solution is 

given by the expression 

xbath = [Zn] / ([Zn] + [Cd]) 

where [Zn] and [Cd] represent the concentrations of Zn and Cd in the liquid solution. On the 

other hand, the relative composition of Zn in the solid ZnxCd1-xS buffer film is given by the 

subindex x, where 

x = stoichiometric composition of Zn in a solid ZnxCd1-xS film 

The relationship between xbath and x is still the subject of research; however, for the purposes of 

reporting results in this document, we make the simplifying assumption that 

x = xbath 

Hence, when in this report we refer to a ZnxCd1-xS film as having a relative fractional zinc 

composition x = 0.3, this statement should be interpreted to mean that the film in question was 

grown using a CBD solution with a concentration characterized by xbath = 0.3. 

2.3.2 Characterizations of ZnxCd1-xS Thin Films 

The chemicals used in this work consist of aqueous solutions of 1.20×10-3 M 

CdCl2·2(1/2)H2O, 1.39×10-3 M NH4Cl, 1.19×10-2 M thiourea (H2NCSNH2), 6.27×10-4 M ZnCl2, 

and 5.27×10-4 M NH3. By varying the relative ratio of Cd and Zn ions, the ZnxCd1-xS films with 

solution composition parameter xbath = 0, 0.3, and 0.5 were deposited on soda-lime-glass (SLG) 

substrates. The bath temperature of the solution was kept around 85 °C. The deposition time 

ranged from 20 to 90 minutes according to the mixture ratio in the solution. The value of pH 

was kept at 7.0 in all sample preparation. SLGs, used as substrates, were cleaned by washing in 

DI water, detergent, ultrasonic cleaner, and 3 step cascade. The cleaned SLG substrates were 

submerged into the starting solution vertically for deposition of ZnCdS films. 

16
2.3.2.1 Thickness of ZnCdS Films 

The film thickness of ZnCdS buffer layers was measured using a Tencor profiler. The 

ZnCdS films were deposited on the surface of soda-lime-glass (SLG) substrates, and were 

covered by #1529 photo resist. The system was then baked for 5 minutes at 110 °C, etched by 

dipping in diluted 10% HCl solution for 20 seconds, rinsed by D.I. water, and finally cleaned by 

acetone and isopropyl alcohol. 

Figure 2-1 shows the film thickness observed as a function of deposition time for ZnCdS 

films with relative zinc compositions x = 0, 0.3, and 0.5. It is important to remark that all the 

films shown in Figure 2-1 were obtained through a single CBD process. For example, for x = 0, 

three SLG substrates were submerged in the CBD solution at the same time, and the glass 

substrates were removed from the bath, one at the time, after 20 min., 30 min, and 40 min, 

respectively. The results show that the growth rate of ZnCdS films with relative zinc 

composition x = 0 (CdS) was faster compared to that of ZnCdS films with x = 0.3, and also the 

growth rate of films with x = 0.3 was faster compared to that of films with x = 0.5 as deposition 

time grew longer. The deposition rate decreases with increasing relative Zn content in the bath. 

0 10 20 30 40 50 60 70 80 90 

Deposition Time (minutes) 

Figure 2-1. Film thickness as a function of deposition time for ZnCdS 

films with relative zinc compositions x = 0, 0.3, and 0.5. 

Figure 2-2 shows the variation of the thickness of the films for different values of relative 

zinc composition x at a constant deposition time of 50 minutes. The thickness of the films is 

found to decrease with an increase of Zn content in the bath. Films showed poor adherence, and 

colors varied from yellow to mostly white with yellow tints, as the composition x increases up to 

0.7. Similar observations have been reported in the literature [2-1]. 

0 

10 

20 

30 

40 

50 

60 

70 

80 

90 

0 

0.3 

0.5 

Film Thickness (nm) 

17
Film Thickness (nm) 

90 

80 

70 

60 

50 

40 

30 

20 

10 

0 

0 0.1 0.2 0.3 0.4 0.5 

[Zn] / ([Zn] + [Cd]) in solution

Figure 2-2. Film thickness variation as a function of relative zinc

composition x for ZnCdS films with deposition time of 50 minutes.

2.3.2.2 Structural Analysis 

A powder X-ray diffractometer with Cu Kα radiation was used to investigate the 

structural properties of ZnCdS films deposited on soda lime glass (SLG) substrates. The ZnCdS 

films prepared by a single CBD process did not show any peaks due to the very thin nature of the 

films. Clear diffraction peaks were obtained using thicker films grown by executing three 

sequential CBD operations, resulting in an approximate thickness of 200 nm (using a bath with 

relative zinc composition x = 0), 100 nm (x = 0.3), and 150 nm (x = 0.5). The resulting powder 

XRD patterns for ZnCdS films with x = 0, 0.3, and 0.5 are shown in Figure 2-3. A comparison 

of the 2θ peak-positions of the JCPDS XRD spectra data for the film with x = 0 (i.e., a CdS film) 

suggests that the as-deposited ZnCdS films have hexagonal (wurtzite group) structures with Xray diffraction peaks corresponding to the (100), (002), (101), (102), (110), (103), (112), and 

(004) planes. The peaks become weaker for films with at the higher Zn compositions x = 0.3, 

and no peaks appear for x = 0.5 reflecting the presence of amorphous structures. 

It is difficult to analyze X-rays of thin films due to their low diffracted intensities 

produced by the film compared to those the substrate and background. A approach for analyzing 

thin films is to use Grazing Incidence-Angle X-ray Diffraction (GIXRD), which intensifies the 

diffracted signal by increasing the path length of the incident X-ray beam by using a very small 

angle (ω) between the incident beam and the film surface. To investigate the structural 

properties of ZnCdS thin layers near the surface or on the surface itself, we used GIXRD with a 

grazing angle ω = 0.4 °. Thin films with relative zinc compositions x = 0, 0.3, and 0.5 were 

deposited by a single CBD process on Mo-coated SLG substrates. The approximate thickness of 

the resulting films was in the range of 50 ~ 60 nm. The GIXRD patterns for ZnCdS films with 

x = 0, 0.3, and 0.5 are shown in Figure 2-4. It is observed that the 2θ diffraction-angle of the 

main peaks shifts to a slightly higher angle with increasing relative Zn composition, due to the 

18
3100 

600 

1100 

1600 

2100 

2600 Intensity (arbitrary units)

350 

550 

750 

950 

1150 Intensity (arbitrary units)

elemental substitutions of Zn for Cd. It is also found that the peaks of the ZnCdS thin film with 

x = 0.5 becomes less pronounced, suggesting a decrease in the crystalline nature of the film at 

that level of zinc content. 

(100)

(101) (002) 

(102)

(110)

(103)

(112)

(004) 

0 

0.3 

0.5

100 

10 20 30 40 50 60 

2 Theta (degree) 

Figure 2-3. Powder XRD patterns of ZnCdS films deposited on the SLG with 

relative zinc composition x = 0 (film thickness d ≈ 200 nm), x = 0.3 (d ≈ 100 nm), 

and x = 0.5 (d ≈ 150 nm). 

1350

Mo (110)

0.5

0.3

0

150 

-50 

2 Theta (degree) 

Figure 2-4. Grazing incidence XRD patterns of ZnCdS films, deposited 

on Mo, with relative zinc compositions x = 0, 0.3, and 0.5.

19 
2.3.2.3 Surface Morphological Analysis 

The surface characteristics of ZnCdS films grown on SLG substrates were examined 

using a field emission scanning electron microscope. The thickness of the ZnCdS films studied 

lied in the range of 70 to 150 nm. Figures 2-5, 2-6, and 2-7 show SEM micrographs of the 

ZnCdS films for different values of relative zinc composition x. The film surface shows larger 

grains as the relative Zn content is increased from x = 0 to x = 0.3. A reduction in crystallinity is 

observed for the film with x = 0.5. This is also consistent with the evidence found in the powder 

and grazing incidence XRD patterns, where much sharper peaks are observed for the films with 

x = 0 and 0.3 compared to the weak peaks observed for the film with x = 0.5. 

Image 1: Actual magnification: 30,000X 

 Accelerating voltage: 15 kV 

Image 2: Selected region

 Actual magnification: 30,000X 

Figure 2-5. SEM micrographs of the surface of a ZnxCd1-xS thin film with 

relative zinc composition x = 0 (CdS). 

Image 3: Actual magnification: 30,000X 

 Accelerating voltage: 1 kV 

Image 4: Selected region

 Actual magnification: 30,000X 

Figure 2-6. SEM micrographs of the surface of a ZnxCd1-xS thin film with 

relative zinc composition x = 0.3. 

20
Image 5: Actual magnification: 30,000X Image 6: Selected region

 Accelerating voltage: 1 kV Actual magnification: 30,000X 

Figure 2-7. SEM micrographs of the surface of a ZnxCd1-xS thin film with 

relative zinc composition x = 0.5. 

2.3.2.4 Optical Analysis 

Optical absorbance and transmittance (T) spectra of ZnCdS films were measured using an 

HP 8453 UV-Visible spectrophotometer over the wavelength range of 300 to 1100 nm. Figure 

2-8 shows the spectral dependence of the transmittance for the ZnCdS films with relative zinc 

composition x = 0.3 deposited on SLG substrates. For very thin ZnCdS films, better than 80% 

transmittance was obtained at short and long wavelengths, while the film with a thickness of 80 

nm has a 70% transmittance for wavelengths longer than 600 nm. The film with a thickness of 

50 nm shows 80% transmittance for wavelengths longer than 600 nm. These values are 

comparable to those reported elsewhere [2-2] [2-3]. 

The thinner films (i.e., those grown with shorter deposition times) have a higher 

transmittance at shorter wavelengths, which is a desirable feature for enhancing the short-circuit 

current. On the other hand, thicker films could assist in avoiding damages caused by the 

subsequent ZnO sputtering deposition process required to build a solar device. Therefore, 

optimization of the buffer layer thickness is essential for improving the overall cell performance. 

Figure 2-9 shows the optical transmittance spectra for ZnCdS films with relative zinc 

composition x = 0 (i.e., a CdS film) and relative zinc composition x = 0.3. The spectral behavior 

of these films shows that the absorption edges shift to shorter wavelengths as the relative Zn 

content increases from x = 0 to 0.3. As indicated in the figure caption, the film thicknesses used 

in the study were 200 nm (x = 0) and 100 nm (x = 0.3). 

The absorption coefficient α (cm-1) for the thin film can be obtained from transmission 

data in the high absorption region through application of the formula 

α = (1/d) × ln(1/T) 

where d (cm) is the thickness of the film, and T is the transmittance of a single-side deposited 

ZnCdS film. Once the absorption coefficient is extracted from the transmission data, the values 

of the optical bang-gap energy (Eg) of ZnCdS films can be estimated using the theoretical 

21
relationship 

αhν = constant × (hν - Eg)

1/2 

where hν represents the photon energy. Figure 2-10 shows the resulting plots of the quantity 

(αhν)

2 versus photon energies in the range 1.0 - 3.5 eV. Finally, from these plots and the 

underlying theoretical relationship it is possible to extract the energy band-gap values reported in 

Table 2-1, which shows that Eg = 2.40 eV for x = 0 (CdS), Eg = 2.55 eV for ZnCdS with x = 0.3, 

and Eg = 2.70 eV for ZnCdS with x = 0.5. Hence, the net effect of including additional Zn 

content in the buffer later is to increase the energy bandgap of film. 

100 

SLG 10 nm 

90 

80 20 nm 

70 50 nm 

80 nm 

60 

50 100 nm 

40 

30 

20 

10 

0 

300 400 500 600 700 800 

Wavelength (nm) 

Figure 2-8. Optical transmittance spectra of ZnCdS films with relative zinc 

composition x = 0.3 for different film thickness. 

Figure 2-9. Optical transmittance spectra for ZnCdS films with relative 

zinc compositions x = 0 (thickness d ≈ 200 nm) and x = 0.3 (d ≈ 100 nm). 

22
0.4 

0.6 

0.8 

1.0 

1.2 

x 

Film Thickness 

(nm) 

Eg 

(eV) 

0 200 2.40 

0.3 100 2.55 

0.5 150 2.70 

1.6 

1.4 

X = 0 4.0 

3.5 

X = 0.3 X = 0.5 1.0 

0.8 

(αhν)2 × (1011cm-2·eV2)

3.0 

(αhν)2 × (1011cm-2·eV2) 2.5 

2.0 

1.5 

(αhν)2 × (1011cm-2·eV2) 0.6 

0.4 

1.0 0.2 

0.2 0.5 

0.0 

0.5 1.0 1.5 2.0 2.5 

hν (eV) 

Figure 2-10. 

3.0 3.5 

0.0 

0.5 1.0 1.5 

Plots of (αhν)

2

 versus 

2.0 2.5 3.0 3.5 

hν (eV) 

0.0 

0.5 1.0 1.5 2.0 2.5 3.0 

hν (eV) 

hν for ZnCdS films with relative zinc 

3.5 

compositions x = 0 (thickness d ≈ 200 nm), x = 0.3 (d ≈ 100 nm), and x = 0.5 (d ≈

150 nm). 

Table 2-1. Estimated energy band gaps of 

ZnCdS films for various relative Zn 

composition values (x). 

2.3.2.5 Electrical Analysis 

The electrical resistivity of the films under study was characterized through I-V 

measurements using an Al/ZnCdS/ITO “sandwich” structure, and also via MMR Hall effect and 

Van der Pauw measurements. 

The I-V measurements yielded the results reported in Table 2-2, which shows the 

resistance, resistivity, and conductivity of ZnCdS thin films as a function of the relative 

composition x. In order to measure the electrical resistivity of CBD-ZnCdS layers, circular 

aluminum contact dots with a thickness of 100 nm and a diameter of 200 μm were deposited by 

electron-beam evaporation. Square glass substrates of 1in × 1in in size were coated with 100 nm 

of indium-tin-oxide (ITO). Prior to the deposition of ZnCdS, the substrates were successively 

cleaned using acetone, methanol, and isopropyl alcohol. The Al/ZnCdS/ITO sandwich structure 

produced good Ohmic contacts for the ZnCdS films. The electrical resistance (R), the electrical 

resistivity (ρ), and electrical conductivity (σ) were estimated using the as-deposited ZnCdS film 

thickness of 70 ~ 80 nm and a contact area of 3.14 × 10-4 cm2

. The measured values of R, ρ, and 

σ are summarized in Table 2-2. The measured electrical resistivities of ZnCdS thin films are in 

the range of 2 ~ 5 × 103 Ω⋅cm, and the resistivity is found to increase with increasing Zn 

composition. 

23
The electrical resistivity was also measured using a four-probe MMR Hall effect and a 

Van der Pauw measurement system. The thickness of the films used in this set of measurements 

lied in the range of 100 to 200 nm. The results indicate that the resistivities of the as-deposited 

ZnCdS films change from approximately 102

 to 7 × 103 Ω·cm as the value of relative zinc 

composition x increases from 0 to 0.5. Theses results show a trend analogous to the I-V results 

obtained using the Al/ZnCdS/ITO sandwich structure. It was also observed that the carrier 

density decreased with increasing Zn content. The carriers found in the films are electrons. 

Similar trends have also been reported elsewhere [2-4] [2-5]. 

Table 2-2. Measured values of electrical resistance (R), resistivity (ρ), and 

conductivity (σ) for as-deposited ZnCdS films on ITO coated glass substrates as a 

function of the relative Zn compositon x. 

x 

Film 

Thickness 

(cm) 

R 

(Ω) 

ρ

(Ω⋅cm) 

σ

(Ω⋅cm)-1 

0 7.0×10-6 53.76 2.41×103

 4.15×10-4 

0.1 7.0×10-6 58.14 2.61×103

 3.83×10-4 

0.3 7.0×10-6 75.19 3.37×103

 2.97×10-4 

0.5 8.0×10-6 117.65 4.62×103

 2.16×10-4 

2.3.3 Photo- J-V and QE of CIGS Solar Cells with ZnCdS Buffer Layers 

Buffer layers of ZnCdS with different Zn content were deposited on CIGS samples 

provided by Energy Photovoltaics Inc. (EPV). The Ga content of the CIGS films was 

approximately 30 ~ 32%, and due to the copper-rich nature of the surface the films were covered 

with a thin chalcogenide Cu2-xSe layer. It is well known that copper-rich CIS and CIGS films 

have high hole densities in the range of 1019 ~ 1020 cm-3 because of the coexistence of a the 

semimetallic Cu2-xSe impurity at the surface. A KCN etching treatment, which consists of 

dipping the as-deposited CIGS film in a KCN solution for several minutes, has been known to be 

an effective treatment to remove Cu chalcogenides [2-6] laying on the surface. Prior to 

depositing the ZnCdS buffer layer, the CIGS films were chemically treated using a 10-wt% KCN 

aqueous solution at room temperature for four minutes. This is followed by deposition of a 

ZnCdS buffer layer via CBD and then by a ZnO window layer via sputtering. A high/low 

resistivity i-ZnO/n+-ZnO bilayer film was then deposited on these samples at EPV. In the next 

step, Ni/Al grid fingers with dimensions of 50 nm/300 nm serving as front contact pads were 

finally deposited by electron-beam evaporation to obtain the finished cells. No antireflective 

(AR) coatings were applied. The measurements taken included the photo- current densityvoltage (photo- J-V) relationships under AM 1.5G (100 mW/cm2

) irradiance at a device 

temperature of 25.0 ± 1°C, and as well as spectral response. 

In the case where CdS is used as the buffer layer, the results given in Table 2-3 shows 

that the CdS/CIGS cells fabricated using a KCN etching treatment have improved the opencircuit voltage (Voc), short-circuit current density (Jsc), and fill factor (F.F.), which yield better 

performance in terms of solar conversion efficiency. The CdS buffer layers were deposited 

using the University of Florida’s baseline CBD. As shown in the table, the CdS/CIGS cell with 

24
ID 21#4, fabricated using a CBD process with a deposition time of 30 minutes and using a KCN 

etching step, realizes the best performance, with Voc = 0.482 V, Jsc = 31.22 mA/cm2, F.F. = 65.10 

%, and Eff. = 9.8 %. 

Figure 2-11 shows spectral response curves for the devices with ID 39#8 and 21#4 of 

Table 2-3, for the purpose of comparing the effect of the KCN etching step. As can be clearly 

seen in the figure, the quantum efficiency of the device utilizing a KCN etching treatment is 

higher than that of the device without the etch treatment. 

Table 2-3. Performance of CdS/CIGS cells with and without a KCN etch 

treatment of the CIGS absorber layers. 

Device 

ID 

Device 

Area 

(cm2

) 

Irradiance 

(mW/cm2

) 

KCN 

Etch 

x Deposition 

Time (min) 

Voc 

(V) 

Jsc 

(mA/cm2

) 

FF 

(%) 

η

(%) 

39#8 0.429 100 No 0 30 0.420 28.04 51.47 6.06 

21#4 0.429 100 Yes 0 30 0.482 31.22 65.10 9.80 

22#1 0.429 100 Yes 0 40 0.487 32.82 59.45 9.51 

300 400 500 600 700 800 900 1000 1100 1200 

Wavelength (nm) 

Figure 2-11. Spectral response curves for the CdS/CIGS solar cells 

shown in Table 2-3 (Device ID: 39#8 and 21#4). 

In the case where ZnCdS is used as the buffer layer, Table 2-4 shows that ZnCdS/CIGS 

cells fabricated using a KCN etching treatment also have improved performance parameters. Quantum Efficiency (%) 

0 

10 

20 

30 

40 

50 

60 

70 

80 

90 

100 

KCN Etch 

No KCN Etch 

25
-10

0

10

20

30

40

-0.2 -0.1 0 0.1 0.2 0.3 0.4 0.5

The effect of the thickness of the ZnCdS buffer layer on the performance of ZnCdS/CIGS 

cells can be assessed from Figure 2-12, which shows photo-J-V characteristics of cells with 

relative zinc composition x = 0.3 and CBD deposition times of 40, 50, and 60 minutes. The 

deposition times were increased from 40 to 60 minutes. The fill factor was found to increase 

when the deposition time reached 60 minutes. Table 2-5 shows the performance parameters of 

these cells. The tested ZnCdS/CIGS cells achieved higher short-circuit currents with respect to 

those of the CdS/ CIGS films reported in Table 2-3, while at the same time realizing comparable 

or slightly higher open-circuit voltages. 

Table 2-4. Performance of ZnCdS/CIGS cells with and without a KCN etch 

treatment of the CIGS absorber layers. 

Device 

ID 

Device 

Area 

(cm2

) 

Irradiance 

(mW/cm2

) 

KCN 

Etch x 

Deposition 

Time 

(min) 

Voc 

(V) 

Jsc 

(mA/cm2

) 

FF 

(%) 

η

(%) 

13#B2 0.100 101.5 No 0.3 45 0.502 30.19 50.68 7.56 

14#A2 0.090 102.1 Yes 0.3 45 0.518 31.24 63.40 10.05 

-0.2 -0.1 0.0 0.1 0.2 0.3 0.4 0.5 

Voltage (V) 

Figure 2-12. Photo-J-V characteristics of ZnCdS/CIGS cells using 

different deposition times and relative zinc composition x = 0.3. The 

CBD-grown ZnCdS buffer layers were fabricated with deposition times of 

40 min (A), 50 min (B), and 60 min (C), in an 85 °C bath. 

-10 

0 

10 

20 

30 

40 

Current Density (mA/cm2 ) 

B 

A 

C 

26
Table 2-5. Performance parameters of ZnCdS/CIGS cells as a function of 

different deposition times for a relative zinc composition x = 0.3 in the buffer 

layer. 

Device 

ID 

Device 

Area 

(cm2

) 

Irradiance 

(mW/cm2

) 

KCN 

Etch 

x Deposition 

Time 

(min) 

Voc 

(V) 

Jsc 

(mA/cm2

) 

FF 

(%) 

η

(%) 

(A)33#3 0.429 100 Yes 0.3 40 0.493 35.16 60.98 10.57 

(B)30#1 0.429 100 Yes 0.3 50 0.487 34.56 60.06 10.11 

(C)32#1 0.429 100 Yes 0.3 60 0.501 34.89 66.81 11.67 

Given that increasing the relative zinc content x can increase the energy band gap of 

ZnCdS, the resulting band gap values are higher than that of CdS. In fact, as discussed in 

Section 1.3.2.4, the ZnCdS film with x = 0.5 had an energy band gap of 2.7 eV, a value that is 

higher than that of CdS at 2.4 eV. Consequently the ZnCdS/CIGS cells are expected to have a 

higher short-circuit current than CdS/CIGS cells because of the higher quantum efficiency in the 

shorter wavelength region. Table 2-6 shows the dependence of the cell performance parameters 

on the relative zinc composition values x = 0.1, 0.2, 0.3, 0.4, and 0.5. The CIGS layers were 

etched by KCN before the CBD buffer layer growth. 

Relevant subsets of the data of Table 2-6 are plotted in Figure 2-13 as an alternative 

visualization of the results. Note from Table 6-1 and Figure 2-13 that the ZnCdS/CIGS cell with 

relative zinc composition x = 0.2 achieved a conversion efficiency of approximately 13% under 

AM1.5G conditions. Also note that starting with the composition x = 0 (corresponding to a 

CdS/CIGS cell) the efficiency increases with increasing relative zinc content in the ZnCdS/CIGS 

cell up to the composition x = 0.2. The data shows then a decrease in efficiency for the 

compositions x = 0.3 and x = 0.4, and finally a marked decrease in efficiency for the zinc-rich 

ZnCdS/CIGS cell with x = 0.5. Given the smaller thickness of the buffer film with composition 

x = 0.2 (cf. Table 2-6), it is reasonable to attribute to that fact the relatively sharp change in 

efficiency between that composition and the efficiency of the cells with x = 0.3 and x = 0.4. It is 

clear, nevertheless, that a zinc content of x = 0.5 is no longer beneficial for the purpose of 

efficiency enhancement. Therefore, the efficiency plot given in Figure 2-13 supports the claim 

that ZnCdS/CIGS cells benefit from the increasing addition of Zn up to a concentration of 

approximately x = 0.3. A reduction in efficiency is realized at the higher Zn content of x = 0.5. 

Figure 2-14 shows the photo-J-V characteristics of ZnCdS/CIGS cells for the cells listed 

in Table 2-6. As mentioned above, the reason for using ZnCdS buffer layers (which have a 

wider band-gap energy than CdS) is the potential for increasing the photo-current generation 

particularly in the spectral region for wavelengths less than 550 nm. 

A comparison of the quantum efficiency of selected cells is shown in Figure 2-15, where 

spectral response curves are given for four ZnCdS/CIGS cells with Zn contents x = 0, 0.1, 0.2, 

and 0.5. At short wavelengths, all the cells featuring ZnCdS as the buffer have better quantum 

efficiency than the cell fabricated with a CdS buffer. Furthermore, at those short wavelengths 

the quantum efficiency of the ZnCdS cells becomes higher as the relative zinc content x 

increases, as can be expected from the widening of the band gap energy caused by the additional 

presence of zinc. This results in an increased value of Jsc, as shown in Table 2-6. The small 

27
difference in cut-off wavelength can be attributed to the difference in Ga content in the CIGS 

absorber layers utilized for fabricating the cells. Further studies and optimization of the ZnCdS 

buffer layer thickness and deposition conditions to improve the film quality are needed for better 

understanding of this buffer layer and achieving the highest performance possible. 

Table 2-6. Performance parameters of ZnCdS/CIGS cells as a function of the 

relative zinc content x of the buffer layer. 

Device 

ID 

Device 

Area 

(cm2

) 

Irradiance 

(mW/cm2

) 

KCN 

Etch 

X Thickness

 (nm) 

Voc 

(V) 

Jsc 

(mA/cm2

) 

FF 

(%) 

η

(%) 

(A)21#4 0.429 100 Yes 0 40 ~ 50 0.482 31.22 65.10 9.80 

(B)38#8 0.429 100 Yes 0.1 60 ~ 70 0.517 32.85 66.31 11.25 

(C)33A#8 0.429 100 Yes 0.2 ~ 40 0.526 35.60 69.53 13.02 

(D)32#1 0.429 100 Yes 0.3 60 ~ 70 0.501 34.89 66.81 11.67 

(E)33B#3 0.429 100 Yes 0.4 60 ~ 70 0.515 34.74 66.85 11.97 

(F)31#1 0.429 100 Yes 0.5 60 ~ 70 0.494 35.17 59.68 10.36 

0.4 

0.5 

0.6 

Voc (V) 

20 

30 

40 

Jsc (mA/cm^2) 

0 0.1 0.2 0.3 0.4 0.5 0 0.1 0.2 0.3 0.4 0.5 

x x 

55 

65 

75 

FF (%) 

9 

10 

11 

12 

13 

14 

Efficiency (%) 

0 0.1 0.2 0.3 0.4 0.5 0 0.1 0.2 0.3 0.4 0.5 

x x 

Figure 2-13. Performance parameters of ZnCdS/CIGS cells as a function of the 

relative zinc content x (graphical representation of the data given in Table 2-6). 

28
Figure 2-14. Photo-J-V characteristics of ZnCdS/CIGS cells with different 

relative zinc contents x = 0 (A), x = 0.1 (B), , x = 0.2 (C) , x = 0.3 (D) , 

x = 0.4 (E), and x = 0.5 (F). 

Figure 2-15. Quantum efficiency comparison between a CdS/CIGS solar 

cell (A), and ZnCdS/CIGS solar cells with relative zinc composition x = 

0.1 (B), x = 0.2 (C), and x = 0.5 (F). 

29
2.4 Conclusions 

We have demonstrated that ZnCdS materials with specific Zn contents have attractive 

physical properties for use as buffer layer materials for the CIGS and CGS solar cells. In 

particular, buffer layers of various relative stoichiometric compositions of Zn, denoted by the 

subindex x in the chemical formula ZnxCd1-xS, were deposited on the SLG substrates. Their 

structural, surface morphology, optical, and electrical properties were characterized and analyzed 

by XRD, SEM, spectrophotometer, and resistivity studies. The experimental studies showed that 

the film growth rate during the CBD process slows down as the zinc content increases; hence, a 

longer deposition time is needed for growing a ZnCdS film than for growing a CdS film to a 

given thickness. The experimental results give an estimate of the expected growth rate as a 

function of the zinc content in the solution. The degree of crystallinity in the film decreased with 

increasing relative zinc content x, leading to an amorphous structure at the concentration level 

x = 0.5. In addition, the resistivity of the ZnCdS films also increased with increasing relative 

zinc content x. Of particular relevance is also the fact that the energy band gap of ZnCdS is 

higher than that of CdS, and that the band gap increases with increasing zinc content x. One 

conclusion is that ZnCdS films with relative zinc composition x = 0.3 showed better than 80% 

transmittance for wavelengths longer than 600 nm with film thicknesses less than 50 nm. 

Because of their high transmittance, ZnCdS can be useful as a buffer layer for the CIGS and 

CGS solar cells. Therefore, it is concluded that from the structural, surface, optical, and 

electrical analyses, the ZnxCd1-xS film with relative Zn content x = 0.3 is an excellent candidate 

as buffer-layer material for the CIGS and CGS solar cells. 

To study the effect of ZnxCd1-xS buffer layers on the performance of CIGS cells, a series 

of buffer films with relative Zn content of x = 0 (CdS), 0.1, 0.2, 0.3, 0.4, and 0.5 were deposited 

on the CIGS samples provided by EPV Inc., and a detailed study of the current-density/voltage 

characteristics and spectral response (QE) was carried out. It was observed that CIGS films 

treated with KCN etching preceding the buffer CBD step lead to improved cell performance for 

CIGS cells deposited with either CdS or ZnCdS buffers. Adding Zn in the ZnCdS buffer layer 

material results in improvement of the performance of ZnCdS/CIGS cells as compared to a Znfree CdS/CIGS cell. In particular, a ZnCdS/CIGS cell with relative zinc composition x = 0.2 

achieved a conversion efficiency of approximately 13 % under AM1.5G conditions, showing 

improved Voc, Jsc, and FF values compared to the CdS/CIGS cell and ZnCdS/CIGS cells with 

different Zn contents. The 13 % efficiency realized by using x = 0.2 in the ZnCdS/CIGS cell 

represents a 30% increase in efficiency over the reference CdS/CIGS solar cell used in this study. 

Furthermore, the quantum efficiency of ZnCdS/CIGS cells becomes higher at shorter 

wavelengths compared to the CdS/CIGS cell as the relative zinc content x increases, until a 

maximum relative zinc composition equal to x = 0.5. 

30
2.5 References Cited 

[2-1] J.M. Dona and J. Herrero, “Chemical Bath Codeposited CdS-ZnS Film Characterization,” 

Thin Solid Films, 268, pp. 5–12 (1995). 

[2-2] T. Yamaguchi, Y. Yamamoto, T. Tanaka, Y. Demizu, and A. Yoshida, “(Cd,Zn)S thin 

Films Prepared by Chemical Bath Deposition for Photovoltaic Devices,” Thin Solid 

Films, 281-282, pp. 375–378 (2003). 

[2-3] A.M. Salem, “Structure, Refractive-index Dispersion and the Optical Absorption Edge of 

Chemically Deposited ZnxCd(1-x)S Thin Films,” Appl. Phys. A 74, pp. 205-211 (1996). 

[2-4] K.T. Ramakrishna Reddy and P. Jayarama Reddy, “Studies of ZnxCd1-xS Films and 

ZnxCd1-xS/CuGaSe2 Heterojunction Solar Cells,” J. Phys. D: Appl. Phys. 25, pp. 1345– 

1348 (1992). 

[2-5] N. Romeo, G. Sberveglieri, and L. Tarricone, “Low-resistivity ZnCdS Films for Use as 

Windows in Heterojunction Solar Cells,” Appl. Phys. Lett. 32, No. 12, pp. 807-809 

(1978). 

[2-6] Y. Hashimoto, N. Kohara, T. Negami, M. Nishitani, and T. Wada, “Surface 

Characterization of Chemically Treated Cu(In, Ga)Se2 Thin Films”, Jpn. J. Appl. Phys. 

Part I, Vol. 35, No. 9A, pp. 4760-4764 (1996). 

31
3 Effects of Buffer Layer Processing on CIGS Thin Films 

Characterized by the Dual-Beam Optical Modulation 

Technique 

3.1 Introduction 

 CuInSe2 (CIS) or CuInGaSe2 (CIGS) thin film solar cells have achieved conversion 

efficiencies as high as 18.8% AM1.5G [3-1]. However, to accelerate the further optimization of 

the device performance, a detailed study and a better understanding of the effects of processing 

on the CIGS/CdS interface and CIGS film transport properties are needed. 

Optimization of CuInSe2 (CIS) or Cu(In,Ga)Se2 (CIGS) thin-film solar cell requires the 

deposition of a thin CdS buffer layer between the CIS absorber layer and the ZnO window layer. 

A great deal of efforts have been devoted to better understanding the role of CdS buffer layers 

deposited by different processing techniques on the performance of CIS and CIGS cells. In order 

to understand why the chemical-bath-deposited (CBD) CdS buffer layer works so well on CIGS 

solar cells, a study of various CdS buffer layer processes on the excess carrier lifetimes and 

material characteristics of CIGS solar cells has been carried out. The CBD grown CdS layer is 

usually thick enough to completely cover the rough surface of CIS or CIGS absorber layers. 

Thus, a pinhole-free conformal coverage buffer layer is formed. Different buffer layer processes 

such as MOCVD and sputtering CdS processes are investigated and compared with CBD 

process. In addition, uncertainties remain as to whether the exchange reactions occurring at the 

interface of CdS and CIGS region with Cd or S ions which chemically passivate the surface, the 

effect of water rinsing the surface of CIGS removing the soluble components [3-2], or 

combinations of these mechanisms affecting the excess carrier lifetimes in the CIGS films. It is 

shown that Cd binds easily on the surface of CIGS film [3-3]. Thus, CdSe and CdInxSey are 

suggested to form and are responsible for graded interface structure [3-2]. The chemical 

modification of the CIGS absorber surface such as Cd- and S-partial electrolyte treatment, DI 

water rinsing treatment and Ar plasma cleaning treatment prior to CdS deposition are also 

investigated to study their effects on excess carrier lifetimes. Both the effects of annealing at 

200oC after the intermediate treatments and MOCVD deposited CdS on as-deposited as well as 

after-intermediate-treated samples on excess carrier lifetimes of the CIGS films are discussed. 

In this study, the reflection mode DBOM technique is employed to investigate the effects 

of different CdS buffer layer processings (i.e., CBD, MOCVD, and sputtering) on the properties 

of CIGS films [3-4]. The results show a significant increase in the DBOM signal (ΔI/I) which is 

related to the change of the free carrier absorption and of the excess carrier lifetime in the 

absorber layer, after the buffer layer processing. Mapping of excess carrier lifetimes across the 

samples can be easily conducted by DBOM technique and it exhibits the substantial 

inhomogeneity of lifetime even for very high quality CIGS samples from various sources. Thus 

the capability of DBOM can be applicable to the quality assessment and process optimization for 

the CIGS films. This novel nondestructive optical technique provides a powerful means for 

assessing the quality of CIGS films, a vital requirement for the viable CIGS solar-cell 

manufacturing. 

32
The reflection mode DBOM technique was applied to study the effects of CdS buffer 

layers deposited by the CBD, MOCVD, and sputtering processes on the excess carrier lifetimes 

in CIGS films. A detailed comparison of the DBOM data versus different processing, the 

measurement and mapping of excess carrier lifetimes in the CIGS films by the DBOM technique 

are presented. Considering the thickness of space charge region between n-type CdS and p-type 

CIGS as zero, the simplified reflection DBOM theory is employed to obtain the effective excess 

carrier lifetimes. A summary of the DBOM measurement results for CIGS films grown by 

National Renewal Energy Laboratory (NREL), Siemens Solar Industries (SSI), and Energy 

Photovoltaics, Inc. (EPV) using different growth techniques, deposited by different buffer layer 

processes, and treated by intermediate surface treatments (i.e., as deposited, DI water rinsed, Cdand S-partial electrolyte treatment, and Ar plasma cleaned) is given in this chapter. 

3.2 Simplified Theory for Reflection-Mode DBOM 

Since CIGS solar cells use thick Mo as backside ohmic contacts, a reflection mode 

DBOM technique has been developed for the measurement of excess carrier lifetimes in the 

CIGS films. The contactless DBOM technique is based on the modulation of the reflected 

intensity of an infrared (IR) probe beam by a visible pump beam which produces free carrier 

absorption in the CIGS film. The fractional change in the reflected intensity of the IR probe 

beam is directly proportional to the effective excess carrier lifetime in the CIGS films. Thus, the 

DBOM technique can be used to determine the excess carrier lifetimes in the CIGS films. It is 

also possible to perform depth profiling of the excess carrier lifetimes in the film by changing the 

wavelength (and hence the absorption depth in the film) of the pump beam. 

Figure 3-1 shows the incident and reflected angles of the pump beam and probe beam in 

the CIGS device structure without p-n junction. The excess electron density, Δn(x), generated by 

the pump beam in the CIGS layer can be obtained by solving the continuity equation for 

electrons and using the boundary conditions at the top (x = 0) and bottom interface (x = tb) of the 

CIGS film, and the result yields 

x t − b x tb 

Δn(x) = αφ0 (1− − R T) cos θ3 (Ce Ln + Ce − L 

−

n − e−αx ) (3-1) Dn (α2 − L

1

2

n

) 1 2 

where C1 = D s )( s D s )( s )e 

(s s )cosh( ) (s s )sinh( ) 

n 1 L n 2 n 2 L

n +

D D 

1 

[t ( )] 

D t D t 

n 

n 

b 1

Ln 

n b n 

2 b 2 

+ − − −

+ + +

− −

( ) ( 

[ ] 

α α α

 (3-2) 

Ln 1 2 Ln 1 2 L2 Ln n 

b (αD + s )( Dn + s ) − (αD − s )( Dn − s )e −[t ( α+ 1

Ln )] 

C = n 1 Ln 2 n 2 Ln 1 (3-3) 2 

2[ Dn (s + s )cosh( tb ) + (s s +

D2

n

)sinh( tb )] Ln 1 2 Ln 1 2 L2

n Ln 

and where θ3 is the reflection angle of the incident pump beam in the CIGS layer, α, φ0, R, and T 

are the absorption coefficient, photon flux, total front surface reflectance, and back-interface 

transmittance at a given pump beam wavelength, while Dn, tb, τ, s1, and s2 are the diffusion 

coefficient, CIGS film thickness, excess carrier lifetime, front-surface recombination velocity 

33
and back-interface recombination velocity in the CIGS films, respectively. 

IR Probe Beam Pump Beam (hν > Eg) 

0 

X j 

t 

b 

t 

f 

R pb R 

pu 

R' 

Tpu Tpb

CdS film 

CIGS film 

θ

θ 2 

θ 1 

Mo film 

x 

3 

τ n 

Figure 3-1. Cross sectional view of CIGS cell used in the simplified DBOM 

theory. 

When pump beam is off, the reflected intensity of the probe beam I can be expressed in 

terms of the transmitted intensity I0, electron density n0, hole density p0, and the optical 

absorption cross section of electrons and holes, σn and σp, respectively, as 

' tb 

I = I R exp[ −2∫ (σ n + σ p )dx] (3-4) 0 n0 p0

0

When pump beam is turned on, the excess carriers generated in the CIGS film will 

modulate the IR reflected beam intensity. Under low injection conditions, the reflected intensity 

of the probe beam I ’ can be express as 

I

' 

= I R exp[ N 0 

' − fc 2Δ σ n t 

n0 

b 

− ∫ 2 

0 

(σ p )dx] + σp 0 (3-5) 

where 

34
tb 

ΔN = ∫ Δn(x)dx (3-6) 

0 

0 ( − − T) cos θ3 − L

tb

n L

tb

n −αtb αφ 1R 1 = 2 [L C (1 n 1 − e ) − L C (1 n 2 − e ) + (e −1)] Dn (α − L

1

2

n 

) α

R’ is the reflectance at Mo/CIGS interface, σfc = σn + σp is the total optical absorption crosssection of electrons and holes. 

 For ΔI<<I, the fractional change in the probe beam intensity ΔI/I (with ΔI = I’ − I) can be 

expressed as 

ΔI ΔI

ln( + ≅ 1) = −2σfc ΔN (3-7) I I 

= −2σ {

αφ0 (1 R− − T) cos θ3 [L C (1 − e 

− L

tb

n ) − L C (1 − e L

tb

n ) +

1 (e−αtb −1)]} fc Dn (α2 − L

1

2 ) n 1 n 2 

α n 

Since the absorption coefficient (α ≈ 8×104

 cm-1) is very high in the CIGS film for the 

wavelength of the pump beam (He-Ne laser, λ=632.8nm), and the bulk trap density in the CIGS 

film is high [3-2][3-5], the following assumptions prevail for DBOM analysis: (i) αLn >> 1, (ii) 

αDn >> s1,s2, (iii) Dn

2

 >> s1s2Ln

2

. Based on these assumptions and solving equations (3-1) 

through (3-7), the fractional change in the probe-beam reflected intensity ΔI/I can be expressed 

as 

ΔI 

= −2σ cosθ φ (1− R T) − τ (1− e−αtb ) (3-8) I fc 30 eff 

where 

τ eff = s1 +

t

t

b

b / τ

(3-9) 

is the effective excess carrier lifetime in the CIGS film, which reduces to the bulk film lifetime if 

the front surface recombination velocity s1 is much smaller than tb/τ. Thus, the excess carrier 

lifetimes in a CIGS film can be obtained by using the values of ΔI and I obtained from the 

DBOM measurements, and the calculated values of R and T. 

The total transmittance and reflectance can be expressed as (∗ denotes the complex 

conjugate) 

T = t × t ∗

 (3-10) 

R = r × r ∗

 (3-11) 

35
The total transmission and reflection coefficients can be expressed as [3-6][3-7] 

−i(δ δ+ ) 2 ttte 1 2 

t = 012 (3-12) δ δ 2 1 + rre −iδ1 + rre −iδ2 + rre −i( 1+ 2 )

01 12 02 

and 

−iδ1 −i( 1+ 2 ) −iδ2 r + re + re 

δ δ + rrre 

r = 0 1 2 012 (3-13) −i δ δ+ 1+ r r e −iδ1 + rre −iδ2 + rre ( 1 2 )

01 12 02 

where 

4π δ = n t θ (3-14) 1 λ CdS CdS cos 2 

4π δ 2 = nCIGS t b cosθ3 (3-15) λ

and where t0, t1, and t2 are the Fresnel transmission coefficients at the air/Cds film, CdS 

film/CIGS film, and CIGS film/Mo film interfaces, respectively; r0, r1, and r2 are the Fresnel 

transmission coefficients at the air/Cds film, CdS film/CIGS film, and CIGS film/Mo film 

interfaces, respectively; δ1 and δ2 are the phase changes that result from the double travel of light 

in the CdS layer and CIGS films; λ is the wavelength of the incident light (laser pump beam); 

tCdS, tb, θ2, and θ3 are the thickness of CdS layer, thickness of CIGS film, and refraction angles in 

the CdS layer and CIGS film, respectively; nCdS is the refraction index of the CdS layer; nCIGS is 

the refractive index of the CIGS film and nCIGS = NCIGS + ik, where NCIGS and k (= αλ/4π) are the 

real and imaginary parts of the refractive index [3-8]; and α is the optical absorption coefficient 

of CIGS film. The Fresnel transmission and reflection coefficients are given by 

t 0 = 2cos θ1 (3-16)

cos θ1 + nCdS cos θ2 

t1 = 2nCdS cos θ2 (3-17)

nCdS cosθ2 + nCIGS cos θ3 

t 2 = 2nCIGS cos θ3 (3-18)

nCIGS cosθ3 + nMo cos θ4 

r0 = cosθ1 − nCdS cos θ2 (3-19)

cosθ1 + nCdS cos θ2 

r = nCdS cos θ2 − nCIGS cos θ3 (3-20) 1 

nCdS cosθ2 + nCIGS cos θ3 

36
r2 = nCIGS cosθ3 − nMo cos θ4 (3-21)

nCIGS cos θ3 + nMo cos θ4 

where nMo is the refraction index of the Mo layer, θ4 is refraction angle in the Mo layer. 

Thus the total transmittance T and reflectance R can be obtained by substituting 

equations (3-12) through (3-21) into equations (3-10) and (3-11). In addition, by measuring the 

total reflectance, the thickness of CdS layer and CIGS film can be determined by using a 

contactless dual beam S-polarized reflectance (DBSPR) technique [3-9], which was originally 

developed to determine the silicon film and buried oxide thickness in SOI materials. Once the 

total transmittance T and reflectance R are obtained from equations (3-10) and (3-11), the excess 

carrier lifetimes can be carried out by substituting T and R into equation (3-8). 

3.3 Experimental Details 

3.3.1 Lifetime Measurement 

As shown in Figure 3-2, a long-pass filter with a cutoff wavelength λc = 1350 nm ,which 

is placed between test samples and detector, is used to block the undesired signal, namely the 

photoluminescence (PL) of CIGS material. A tungsten lamp and a He-Ne CW laser (λ = 

632.8nm) are used as probe beam and pump beam, respectively. 

Tungsten 

Lamp Pinhole 

Lens Sample 

Lens IR Filter 

Under 

Test 

Ge 

He-Ne 

Laser Lock-in 

Chopper 

Planar 

Mirror 

Chopper 

Laser 

Power 

Meter 

Detector 

amplifier 

line 

filter 

Long 

Pass 

Filter 

Figure 3-2. Schematic diagram of reflection mode DBOM setup with a long pass 

filter. 

37
The experimental procedure is composed of two steps. First, the reflected intensity I of 

an IR probe beam (with hν < Eg of CIGS) is measured by chopping the tungsten lamp, and the 

signal is detected with a lock-in amplifier. Second, the change in the reflected beam intensity ΔI, 

due to a modulated pump beam (hν ≥ Eg), is measured by using the chopped pump beam, and the 

resulting IR reflected beam intensity is measured using a lock-in amplifier. 

3.3.2 Buffer Layer Processing 

Three different CdS processing methods (CBD, MOCVD, and sputtering) were used for 

the deposition of CdS buffer layers on the Mo-coated CIGS samples, which were supplied by 

National Renewal Energy Laboratory (NREL), Siemens Solar Industries (SSI), and Energy 

Photovoltaics Inc. (EPV). Prior to buffer layer processing, intermediate surface treatments (i.e., 

as deposited, DI water rinsed, Cd- and S-partial electrolyte treatment, and Ar plasma cleaned) 

were also applied to some of these samples. The DBOM measurements were performed on 

samples with or without the CdS buffer layers deposited by these processing techniques. 

Two sets of CIGS films were grown on soda-lime glass substrates coated with thick Mo 

films at NREL. For set-one samples, EPMA data taken on these substrates showed Cu 23.72% 

(± 0.123%), In 18.76 (± 0.11%), Ga 7.2% (± 0.06%), and Se 50.3% (± 0.09%). All of these 

substrates were further divided, and ESCA measurements of the surface and near-surface 

composition profiles were taken. The ESCA data showed that surfaces of all these films were 

significantly enriched with indium with respect to their average compositions and the 

composition in the copper-depleted near-surface region. Figure 3-3 schematically illustrates four 

processing conditions studied and measurements conducted in each case. The baseline CBD 

solution consisted of NH3 (1.44×10-4M), NH4Cl (7.43×10-4M), CdCl2 (2.4×10-4M), and 

NH2CSNH2 (2.4×10-3M). The cadmium partial electrolyte consisted of a solution with all the 

same reactant except thiourea, and the sulfur partial electrolyte contained all the baseline 

reactants in the same concentration except the cadmium salt. The aqueous processing was 

conducted at a bath temperature of 80-85°C and was continuously stirred. The processing time 

for the partial electrolyte cases is identical to the baseline CBD process, which produced a 

uniform, compact, coherent, and colloid-free film with a thickness of about 800Å. MOCVD was 

performed in a horizontal atmospheric pressure cold wall reactor using dimethyl cadmium and 

hydrogen sulfide reactants in hydrogen at a substrate temperature of 160±10 °C and a VI/II 

reactant ratio of 100/1 for 2.8 minutes. The MOCVD growths were prenucleated with 

approximately one minute of exposure to H2S before the flow of dimethyl cadmium began. The 

DBOM measurements were conducted on sixteen distinct locations for each substrate at each 

step in the process per Fig. 3-3. 

38
 CBD CdS

Sample A 200o

C annealing

 2 minutes

 DBOM

 DBOM

 DBOM

 DBOM

 DBOM

 DBOM

 DBOM

 DBOM

 DBOM

 DBOM

 MOCVD CdS

Sample B Cd-partial 200o

C annealing

electrolyte 2 minutes

 MOCVD CdS S-partial Sample C 200o

C annealing electrolyte 2 minutes 

Sample D MOCVD CdS 

Figure 3-3. Process flow sequence of NREL CIGS samples for the four 

experimental conditions. 

For set-two NREL CIGS samples, the CdS buffer layer was deposited by using the CBD 

(NREL), sputtering (MRG), and MOCVD (University of Florida) techniques. The processing 

conditions studied and measurements conducted for these samples are shown in Figure 3-4. 

These samples measured by DBOM technique were not fabricated into devices but are believed 

to be comparable to the similar samples with efficiencies of around 14%. 

Sample number M1517 Sample number M1518 

 As deposited DI water Cd electrolyte 

C B A CBD CdS 

MOCVD CdS 

Sputtering CdS 

 As deposited DI water Cd electrolyte 

C B A

 As deposited

F 

 DI water

 E 

 Cd electrolyte 

D

 As deposited DI water Cd electrolyte 

I H G 

Figure 3-4. The twelve experimental conditions for NREL CIGS samples. 

39
In addition to the samples fabricated by NREL, CIGS films were provided from two 

different sources (SSI and EPV). The CdS buffer layers for these samples were deposited by 

using either the CBD (University of Florida) or sputtering (MRG) technique. The CdS 

deposition (CBD and sputtering) was carried out on the as-deposited CIGS samples and on CIGS 

samples treated with DI water rinsing and Ar plasma cleaning. These samples measured by 

DBOM technique were not fabricated into solar cell but are believed to be comparable to similar 

samples with efficiencies of around 3-5% (best one 9%) fabricated by NREL. 

3.4 Results and Discussion 

The following results are obtained by using the simplified model (no p-n junction) and 

the data of ΔI and I measured by the reflection-mode DBOM technique. 

3.4.1 Measurements of NREL Samples 

The different processing conditions and measurements conducted in each case are 

illustrated in the Fig. 3-3 for the first set of NREL samples. Fig. 3-5 is a box plot showing a 

statistically summary of the excess carrier lifetime data taken for each experiment case, initially 

(as-received), after intermediate surface treatments (Cd- and S-partial electrolyte), and after the 

completion of processing. Several results are evident. First, the initial lifetimes in the CIGS 

films, in the range of smaller than 3ns, are typically improved by a factor of two to three in every 

case except the aqueous sulfur electrolyte exposure. This result is strong negative evidence for 

attributing the beneficial effects on CIGS device performance of the CBD CdS process to sulfur 

passivation, an effect well established in the case of III-V compound semiconductors. However, 

the data does not support the alternative hypothesis that cadmium ion reactions with the surface 

during the CBD process alone are responsible for the dramatic lifetime improvement shown in 

the cadmium partial electrolyte case subsequent to the MOCVD growth of CdS, because the 

lifetimes measured immediately after aqueous processing showed no significant change 

compared to their initial values. Indeed, considering the substantially poorer initial lifetime in 

the MOCVD control sample compared to the CBD control and cadmium partial electrolyte 

samples, it appears that the beneficial effect is provided equally well by the MOCVD growth 

process alone. Since all the samples were baked in air before final measurement except the 

MOCVD control sample, the possibility that oxidation or higher temperature interdiffusion 

effects are required to achieve high lifetimes is not supported by this data. Since the MOCVD 

control (which was not subjected to any aqueous processing) had a beneficial effect on lifetimes, 

and the aqueous sulfur electrolyte exposure a neutral effect, water exposure alone cannot be 

responsible for the observed lifetime increases. 

40
9 

8 

7 

6 

5 

4 

3 

C d

2 

1 

(A ) (B ) (C ) 

0 

 e le c tro lyte 

S e le c tro lyte 

CBD 

M O C V D 

M O C V D 

M O C V D 

(D ) 

(B ) 

(C ) 

(C ) 

(A ) (B ) 

(D ) 

A S d e p o s ite d A fte r tre a tm e n ts A fte r C d S d e p o s itio n 

Figure 3-5. A statistical summary of excess carrier lifetimes for the first set 

NREL CIGS samples. 

Figures 3-6, 3-7, 3-8, and 3-9 are the contour-line plots of DBOM data showing the 

spatial inhomogeneity in excess carrier lifetimes exhibited in these very high quality films 

initially, after surface treatments, and even for the CIGS samples after the CBD and MOCVD 

processings. Future studies will use this capability in conjunction with the materials 

characterization to obtain a correlation with the material parameters and defect structures, which 

control the recombination processes in these device structures. 

41
15 

10 

5 

0 

1 1

1.2 1.2 

1.4 

1.4 

1.5 1.5 

1.2 1.2 

1 1 0.8 

15

5 

10 

0

7 7

8 

9 

6 6 

5 

5 

4 

0 4 8 12 0 4 8 12 

X-position [mm] X-position [mm] 

(a) Initially (b) After CBD deposited CdS 

Figure 3-6. Contour line plot of excess carrier lifetimes for NREL CIGS sample 

(a) initially (as deposited) (b) after CBD deposited CdS. 

42
0 

5 

10 

15 1 

1.5 

1.5 

2 

0 4 8 12 

X-position [mm] 

(a) Initially 

15 

10

5 

0 

15 

5 

10 

0 

1 

1.5 

2 

0.5 

0.75 

1.75 4 

5 

6 

6 

7 

7 

8 

9 

10 

9 

8 

0 4 8 12 0 4 8 12 

X-position [mm] X-position [mm] 

(b) After surface treatment (c) After MOCVD deposited CdS 

Figure 3-7. Contour line plots of excess carrier lifetimes for NREL CIGS sample 

(a) initially (as deposited) (b) after surface treatment (Cd-partial electrolyte) (c) 

after MOCVD deposited CdS. 

43
0 

5 

10 

15 

2 

2 

2 

2.25 

1.5 

1 

0.5 

0 4 8 12 

X-position [mm] 

(a) Initially 

15 

10

5 

0 

15 

5 

10 

0 

2 

2 

2 

1.75 

1.75 1 1.5 

0.5 

1.8 

2 

2.4 

2.2 

2.6 

2.8 

3 

0 4 8 12 0 4 8 12 

X-position [mm] X-position [mm] 

(b) After surface treatment (c) After MOCVD deposited CdS 

Figure 3-8. Contour line plots of excess carrier lifetimes for NREL CIGS sample 

(a) initially (as deposited) (b) after surface treatment (S-partial electrolyte) (c) 

after MOCVD deposited CdS. 

44
15 15

10 

0 4 8 12 

1.5 

1 

0.5 

0.5 

0.26 

0.26 

10 

0 4 8 12 

2.5 

3 

4 

5 

5 

6 

7 

8 

5 5

0 0

X-position [mm] X-position [mm] 

(a) Initially (b) After MOCVD deposited CdS 

Figure 3-9. Contour line plot of excess carrier lifetimes for NREL CIGS sample 

(a) initially (as deposited) (b) after MOCVD deposited CdS. 

It is known that CIGS is the material with photoluminescence (PL). The 

photoluminescence is the undesired signal for the DBOM measurements. After both the 

deposition of CdS and annealing process on the sample D, the signal ΔI of DBOM measurement 

can not be measured correctly. Generally the signal ΔI detected by a lock-in amplifier is in the 

range of microvolts (μV). Thus, it is possible that the photoluminescence emission from the 

CIGS sample affects the DBOM measurement. It is shown that the higher temperature of 

annealing process on CIGS samples results in the higher PL intensity emission from CIGS 

samples [3-10]. This is the reason why the DBOM measurement can not be performed correctly 

on the sample D after annealing process. According to the PL spectra of CIGS samples [3-10][3­

11], the peak positions of PL intensity are around at wavelength λ = 1280 nm. Therefore a longpass filter with a cutoff wavelength λc = 1350 nm is used to block the undesired signal, 

photoluminescence emission from CIGS material, and it is set between test samples and detector 

as shown in Fig. 3-2. With the long-pass filter, the reflection mode DBOM measurement can be 

performed on CIGS samples without a hitch. 

The results for the second set of NREL CIGS samples were determined by the DBOM 

technique, and are summarized in Tables 3-1, 3-2, and 3-3, Figures 3-10, 3-11, and 3-12. Table 

3-1 and Fig. 3-10 show a statistical summary of the excess carrier lifetimes taken for the 

experimental cases of as deposited, after DI water rinsed, after Cd partial electrolyte, and after 

annealing at 200o

C for 2 minutes for the samples without CdS. Table 3-2 and Fig. 3-11 

summarize the excess carrier lifetimes of the samples without CdS and with CBD, MOCVD, and 

sputtering deposited CdS and with different surface treatments. For discussing the MOCVD CdS 

process, Table 3-3 and Fig. 3-12 show the excess carrier lifetimes for each experimental 

condition: as deposited, after surface treatments, and after MOCVD CdS. And the contour-line 

plots of excess carrier lifetimes for three CIGS samples as deposited initially, after surface 

treatments, and MOCVD CdS deposition are shown in Figures 3-13, 3-14, and 3-15. It should be 

pointed out that there is a considerable variation in the excess carrier lifetimes for the as45
deposited CIGS samples without CdS, as is shown in Table 3-1 and Fig. 3-10, respectively. One 

of the difficulties in interpreting the DBOM results in these CIGS samples is the fact that we 

have only the samples with the CBD as well as sputtering CdS deposition and we do not have the 

DBOM measurements (and hence the excess carrier lifetimes) on these CIGS samples prior the 

deposition of CdS. 

Table 3-1. The excess carrier lifetimes τeff (ns) of the CIGS samples without 

CdS. 

sample # As deposited DI water rinsed Cd partial electrolyte Annealing 

M1517-A 1.46 --- 1.34 1.43 

M1517-B 1.90 1.73 --- 1.75 

M1517-C 0.77 --- --- --- 

Table 3-2. The average excess carrier lifetimes of NREL samples. 

As deposited DI water rinsed Cd partial electrolyte

 sample # τeff (ns) sample # τeff (ns) sample # τeff (ns) 

no CdS M1517C 0.77 M1517B 1.75 M1517A 1.43 

CBD CdS M1518C 3.95 M1518B 2.76 M1518A 3.56 

MOCVD CdS M1518F 17.6 M1518E 1.91 M1518D 4.67 

Sputtered CdS M1518I 12.7 M1518H 7.64 M1518G 5.48 

Table 3-3. The excess carrier lifetimes τeff (ns) of NREL CIGS samples for 

MOCVD CdS. 

sample # As deposited DI water rinsed Cd partial electrolyte MOVCD CdS 

M1518-D 2.29 --- 2.08 4.67 

M1518-E 1.71 1.70 --- 1.91 

M1518-F 0.53 --- --- 17.6 

46
2. 

0. 

0 

1 

1. 

2 

DI 

Cd 

(A) 

(B) 

(B) (B) 

(C) (C) (C) 

Cd electrolyte 

2.5 

0.5 

As deposited After surface treatment After annealing 

Figure 3-10. A statistical summary of excess carrier lifetimes for the second set 

NREL CIGS samples (M1517-A,B and C), the experimental conditions: as 

deposited, after surface treatments and after annealing. 

47
0 

2 

4 

6 

8 

10 

12 

14 

16 

no CdS 

no 

CdS no 

CdS 

CBD 

CBD 

CBD 

MOCVD 

MOCVD 

MOCVD 

sputtered 

sputtered 

sputtered 

As deposited DI water rinsed Cd-partial electrolyte 

0

2

4

6

8

10

12

14

16

no CdS

no

CdS no

CdS

CBD

CBD

CBD

MOCVD

MOCVD

MOCVD

sputtered

sputtered

sputtered

As deposited DI water rinsed Cd-partial electrolyte

Figure 3-11. A statistical summary of excess carrier lifetime for the second set 

NREL CIGS samples without CdS and with CBD, MOCVD, and sputtering 

deposited CdS and with surface treatments. 

48
 

12 

10 

8 

6 

4 

2 

0 

(D) 

(D) 

(E) 

DI water 

(E) 

(F) Cd electrolyte 

(F) 

(E) 

(F) 

As deposited After treatment After MOCVD CdS 

Figure 3-12. A statistical summary of excess carrier lifetimes for MOCVD CdS 

deposited CIGS samples. 

49
12 

6 

0 

0 6 12 

2.4 

2.45 

2.3 2. 

2.2 2.2 

X-position [mm] 

(a) Initially 

12 

6 6 

0 6 12 

2.1 

2.1 

2.15 

2.05 

2 

12 

0 0 

X-position [mm] 0 6 12 

4

4 

3

5 

6

4.5 

X-position [mm] 

(b) After surface treatment (c) After MOCVD deposited CdS 

Figure 3-13. Contour line plots of excess carrier lifetimes for NREL CIGS 

sample (a) initially (as deposited) (b) after surface treatment (Cd-partial 

electrolyte) (c) after MOCVD deposited CdS. 

50
0 

6 

12 

1.9

1.8 

1.6 

1.5 

1.4 

1.7 

9

0 6 12 

X-position [mm] 

(a) Initially 

6 

12 12 

1.8 

2 

1.6 

1.4 

1.3 

6 

0 0 

2 

5 

1.5 

4 

3 

0 6 12 0 6 12 

X-position [mm] X-position [mm] 

(b) After surface treatment (c) After MOCVD deposited CdS 

Figure 3-14. Contour line plots of excess carrier lifetimes for NREL CIGS 

sample (a) initially (as deposited) (b) after surface treatment (DI water rinsed) (c) 

after MOCVD deposited CdS. 

51
12 

6 

0

1 6 12 

0.4 

0.6 

0.8 

1 

0.3 

12 

6 

0 

0 6 12 

12 

11 

11 

10 

9 

8 

7 

6 

5 

X-position [mm] X-position [mm] 

(a) Initially (b) After MOCVD deposited CdS 

Figure 3-15. Contour line plot of excess carrier lifetimes for NREL CIGS sample 

(a) initially (as deposited) (b) after MOCVD deposited CdS. 

The effect of CdS deposition on the excess carrier lifetimes in CIGS films is obvious. 

The results reveal that the excess carrier lifetimes are significantly increased with CdS 

deposition. The CIGS samples with the sputtered CdS show a dramatic increase in excess carrier 

lifetimes for both the as-deposited and after surface treatments. However, no clear evidence of 

positive effect of the surface treatments (DI water rinsing and Cd partial electrolyte treatments) 

and annealing on the excess carrier lifetimes in these CIGS samples is obtained from this study. 

It is interesting to note that the most dramatic increase in the excess carrier lifetimes was 

observed in the sputtered CdS deposited samples with no intermediate surface treatments. Also, 

the direct processing of sputtered CdS on the as-grown CIGS sample yields the longest excess 

carrier lifetime when compared to the CdS deposited CIGS samples. 

3.4.2 Measurements of SSI and EPV Samples 

The DBOM results for these samples are shown in Fig. 3-16. No detectable DBOM 

signal was obtained on all CIGS samples without the CdS buffer layers, as shown in Fig. 3-16. It 

is obvious that the CdS buffer layers increase ΔI/I values (and hence excess carrier lifetimes) for 

both the DI water-treated and as-deposited samples. This is especially true for the sputtering 

CdS deposited CIGS samples which show a dramatic increase of the DBOM signal for those DI 

water-treated and as deposited samples, as illustrated in Fig. 3-16. For both the CBD and 

sputtering CdS deposited samples, the DI water rinsing treatment decreases the excess carrier 

lifetimes. It is interesting to point out that the Ar plasma cleaning process dramatically increases 

the DBOM signal for the samples without the CdS buffer layer. The Ar plasma cleaning has the 

effects of Ar ion bombardment on the surface of the CIGS samples. Thus a smaller depletion 

width could have been created in the surface junction because of the Ar-plasma cleaning process. 

After Ar-plasma cleaning process, according to the XPS analysis the data show a clear evidence 

of chemical change in the samples, especially near the surface. The Cu/In ratio is restored to 

around unity although the starting surface composition was heavily In-rich. The depletion width 

52
CBD 

CBD 

CBD 

CBD 

(B) 

(A) 

sputtered 

sputtered 

sputtered 

(B) 

(A) 

(B) 

sputtered 

(A) 

(A) 

(B) 

no CdS no CdS 

without 

(B) 

DI 

water 

(A) 

(B) 

(A) 

DI water 

(A) 

(B) 

DI water 

sputtered 

(A) 

(B) 

DI water 

+ + 

CBD

CBD

CB

CBD

(B)

(A)

sputtered

sputtered

sputtered

(B)

(A)

(B)

sputtered

(A)

(A)

(B)

no CdS no CdS

without 

(B)

DI

water

(A)

(B)

(A)

DI water

(A)

(B)

DI water

sputtered 

(A)

(B)

DI water

+ +

created due to the chemical change near the surface of the CIGS film could be the reason why 

the similar samples fabricated into devices have an increase in Voc compared to the cells 

fabricated by using the as-grown and DI water rinsed samples. The created small depletion 

width and the band bending do have an effect on the results of DBOM measurements, which 

show an increase in the DBOM signal. However, after the sputtering CdS process on the CIGS 

samples the DBOM signal decreases. 

70

As DI water DI water Aplas A plas

(A):SSI 

(B):EPV 60

50

40

30

20

10

0

Figure 3-16. A statistical summary of excess carrier lifetimes for SSI and EPV 

CIGS samples without CdS and with CBD and sputtering deposited CdS, and 

with surface treatments. 

3.5 Conclusions 

We have demonstrated a contactless reflection mode DBOM technique developed 

specifically for determining the excess carrier lifetimes in the CIS and CIGS films. A study of 

the effects of CBD, MOCVD and sputtering deposited CdS on the excess carrier lifetimes in the 

CIGS films has been conducted. The results of this DBOM study for the CIGS samples from 

different laboratories (NREL, EPV, and SSI) are summarized as follows: the CBD, MOCVD and 

sputtering deposited CdS buffer-layer formation increase the excess carrier lifetimes for all asdeposited, water-rinsed, S- and Cd-partial electrolyte-treated samples. The CdS deposition 

increases ΔI/I less for water-rinsed, S- and Cd-partial electrolyte-treated samples than for asdeposited samples. Sputtering CdS deposition decreases the excess carrier lifetimes for the Ar53
plasma cleaned samples. However, prior to the CdS deposition, Ar-plasma cleaning process 

exclusively increases the excess carrier lifetimes of CIGS films. The intermediate surface 

treatments (i.e., DI water rinsed, Cd-partial electrolyte treatment, and Ar plasma cleaned) are not 

a positive effect on excess carrier lifetimes for CdS deposited CIGS samples. Especially, the 

sputtering of CdS dramatically increases the DBOM signal ΔI/I for as-deposited, water- and Cdpartial electrolyte-treated samples. The relative free carrier absorption change as measured by 

ΔI/I shows consistent trends for the CIGS films from several laboratories fabricated by the 

different techniques. The free carrier absorption change measured by the DBOM signal can be 

directly related to the quality of CIGS films studied in this work. 

Further study of the correlation between the excess carrier lifetimes and the performance 

parameters of CIGS solar cells will be carried out in order to optimize the processing and growth 

conditions of CIGS films. 

3.6 References cited 

 [3-1] M. Contreras, B. Egass, K. Ramanathan, J. Hiltner, A. Swartzland, F. Hasoon, and R. 

Noufi, Progress in Photovoltaics, Vol. 7, pp. 311, 1999. 

[3-2] D. Lincot, R. Ortega-Borges, J. Vedel, M. Ruckh, J. Kessler, K. O. Velthaus, D. Hariskos, 

and H. W. Schock, 11th E. C. Photovoltaic Solar Energy Conference, p. 872, 1992. 

[3-3] T. M. Friedlmeier, D. Braunger, D. Hariskos, M. Kaiser, H. N. Wanka, and H. W. Schock, 

Conference Record of the 25th IEEE PVSC, p. 845, 1996. 

[3-4] Sheng S. Li, B. J. Stanbery, C. H. Huang, C. H. Chang, T. J. Anderson, and Y. S. Chang, 

Conference Record of the 25th IEEE Photovoltaic Specialists Conference, p. 821, 1996. 

[3-5] Y.-S. Chang, Sheng S. Li, R. Datta, and L. P. Allen, Solid-State Electronics, Vol. 41, p. 

1189, 1997. 

[3-6] L. Ward, The Optical Constants of Bulk Materials and Films, Adam Hilger, England, 

1988. 

[3-7] M. Levy, E. Scheild, S. Cristoloveanu, and P. Hemment, Thin Solid Films, Vol. 148, p. 

127, 1987. 

[3-8] J. I. Pankove, Optical Process in Semiconductors, Dover, New York, 1975. 

[3-9] Yun-Shan Chang and Sheng S. Li, Solid-State Electronics, 38(2), p. 297, 1995. 

[3-10] K. Ramanathan, M. A. Contreras, J. R. Tuttil, J. Keane, J. Webb, S. Asher, D. Niles, R. 

Dhere, A. L. Tennant, F. S. Hasoon, and R. Noufi, Conference Record of the 25th IEEE 

Photovoltaic Specialists Conference, p. 837, 1996. 

[3-11] S. Zott and K. Leo, Conference Record of the 25th IEEE Photovoltaic Specialists 

Conference, p. 817, 1996. 

54
4 Epitaxial Growth of CuInSe2 at High Temperature 

4.1 Introduction 

It was possible to grow CuInSe2 (CIS) epitaxially at a relatively low temperature (360 

°C) for Cu-rich compositions. The explanation is that the excess Cu-Se phase acts as a surfactant 

to increase the mobility of surface adatoms and thus allow the adatoms to approach the 

nucleation sites and to grow epitaxially. It has also been suggested that a Cu-Se liquid phase 

exists and enhances the epitaxial growth of CIS, although no direct evidence has been reported. 

For the In-rich composition, the CIS film was not grown epitaxially and it showed only the 

domains with larger grain size. The polycrystalline growth was attributed to insufficient Cu to 

form the Cu-Se secondary phase, which increases the surface mobility and thus enhances twodimensional-layer growth. Therefore, it was postulated that at higher growth temperature, the 

surface mobility, which increases exponentially with temperature, will be sufficient for epitaxial 

growth of In-rich films without the help of the Cu-Se phase. For these reasons, the growth of CIS 

films with both Cu-rich and In-rich compositions was performed at a higher growth temperature 

(464 vs. 360 °C). 

Furthermore, the role of Na on the growth of CIS films is not well understood. From the 

literature review, it has been suggested that Na affects the growth process to increase the grain 

size by acting as a surfactant. Therefore, it is hypothesized that Na will also enhance epitaxial 

growth of CIS film, even without the help of the Cu-Se secondary phase. Four CIS films were 

grown on (100) GaAs, Cu-rich and In-rich, each with and without Na dosing. The growth mode 

of the films was studied by several techniques including ICP-OES, SEM, XRD, and TEM. 

4.2 Experimental Approach 

The three-temperature sequence used for the epitaxial growth of CIS at low growth 

temperature was tried using higher initial and final temperatures, and resulted in polycrystalline 

CIS or poor crystalline quality. This is attributed to a lower density of nucleation sites with 

higher initial temperature to give a longer migration length of adatoms to reach the nucleation 

site and achieve epitaxial growth. 

Four CIS films were grown at the conditions summarized below. A constant growth 

temperature of 464 °C was used for the growth of CIS films. Because of a limitation in the 

system design and an increase in the background pressure of the reactor during the growth, this 

temperature was the maximum achievable one in the system. The background pressure during 

the deposition process was in the range of 2×10-7 to 6×10-7 Torr. The direction of the rotation 

(counter-clockwise) and the rotating speed (12 rpm) were the same as in low temperature growth 

experiments. Again, semi-insulating (100) GaAs wafers were used as substrates. The bulk 

composition of each CIS film was measured by ICP-OES for companion films grown on 

molybdenum-coated glass substrate by dissolving the film in 15% HNO3 solution. Two CIS 

films, one with Cu-rich composition (W558) and the other with In-rich composition (W560) 

were grown. NaF was used as the Na precursor and delivered to the film from a Knudsen cell. 

Both Cu-rich (W584) and In-rich (W590) CIS films were dosed with Na with a Na flux 

impinging on the surface continuously and initiated at the same time as film growth. 

55
4.3 Composition, Phase and Surface Morphology of CIS Films 

Table 4-1 summarizes the bulk composition of CIS films. Sample W558 is Cu-rich 

(Cu/In = 1.10), while sample W560 is slightly In-rich (Cu/In = 0.99). 

The level of Na dosing lies between 0.5 and 1 %, as stated in Table 4-1. CIGS absorber 

layers for high efficiency cells typically show 0.1 at.% of Na inclusion [4-1]. As such a small 

amount of Na in the films is difficult to detect, secondary ion mass spectroscopy (SIMS) was 

used to verify the Na level in the film. Depth profiling was also performed to investigate the 

composition of each element along the depth of CIS film. SIMS analysis was performed using a 

Perkin-Elmer PHI 6600 SIMS system operating with a 5 keV Cs primary ion beam. The depth 

profile of the composition of each element is shown in Figure 4-1. In this figure, the depth was 

calculated from the value of the crater depth after SIMS analysis (1396 nm). The vertical dotted 

line represents the CIS film/GaAs wafer interface and was estimated by the film thickness of 550 

nm as observed in Figure 4-11b by SEM. The element profiles shown in Figure 4-1 indicate 

considerable diffusion of Cu, In, and Se into the GaAs and the diffusion of Ga from the GaAs 

into the CIS film. This is, however, believed to be an artifact from uneven sputtering of the rough 

film surface (AFM image shown in Figure 4-5 with estimated RMS roughness of 3.25 nm). In 

addition, the raster size was large (400×400 μm). This conclusion is supported from a TEMEDX scan, which shows that In and Se were not detected in the GaAs substrate (detection limit ~ 

0.5%) and Ga is not detected in CIS film though it was not plotted. However, Na was observed 

to diffuse into the GaAs substrate by TEM-EDX scan. The diffusion of Cu is also supported 

from the observation of PL peak related to the CuGa - Nai defect states in GaAs. 

Table 4-1. Bulk composition of CIS films on Mo-coated glass determined by ICP. 

Sample Cu (at. %) In (at. %) Se (at.%) (Δs) Na (at. %) Cu/In 

(Δm) 

Growth rate 

(Å/s) 

W558 26.76 24.38 48.87 (-0.022) - 1.10 (0.098) 0.59 

W560 25.30 25.58 49.12 (-0.037) - 0.99 (-0.011) 0.48 

W584 27.06 24.75 47.69 (-0.059) 0.50 1.09 (0.093) 0.52 

W590 23.60 26.17 49.24 (-0.036) 0.99 0.90 (-0.098) 0.54 

56
Figure 4-1. Element depth profiles of the In-rich, Na-dosed CIS sample (W590). 

To investigate the effect of the Cu/In ratio and inclusion of Na dosing on the crystal 

quality and growth mode of CIS films at this higher growth temperature, the surface 

morphologies of the films were investigated by SEM as shown in Figures 4-2 through 4-4. 

Figure 4-5 is an AFM image as W590 has a very smooth surface that does not show the contrast 

in a SEM image. The Cu-rich CIS films without (W558, Figure 4-2) and with Na dosing (W584, 

Figure 4-4) show the island structures embedded in the matrix region similar to the Cu-rich CIS 

films grown at low temperature. These island structures in Cu-rich CIS films show elongated and 

faceted structures oriented in [0 1 -1] direction. The density of islands in these two Cu-rich 

samples is ~ 5×105

 cm-2 for both samples (W558 and W584) and lower than the density of 

islands in Cu-rich CIS sample (1.1×106

 cm-2) with Cu/In ratio of 1.09. 

2 μm 

20 μm 

[0 1 -1] 

(a) (b) 

Figure 4-2. SEM image of the surface of the Cu-rich CIS film (W558, Cu/In = 1.10): 

(a) 500X, (b) island region at 10,000X. 

57
20 μm 

2 μm [0 1 -1] 

(a) (b) 

Figure 4-3. SEM image of the surface of the In-rich CIS film (W560, Cu/In = 0.99): 

(a) 500X, (b) island region at 10,000X. 

There are a few grains boundaries observed in the island regions of the Cu-rich CIS film 

without Na dosing (W558) while the island regions in the Cu-rich CIS film dosed with Na 

(W584) do not show faceted grain boundaries. 

20 μm [0 1 -1] 2 μm

(a) (b) 

Figure 4-4. SEM image of the surface of the Cu-rich, Na dosed CIS film (W584, Cu/In = 1.10):

(a) 1,000X, (b) island region at 10,000X. 

58
Z = 20 nm/div Z = 20 nm/div 

[0 1 -1] [0 1 -1] 

(a) (b) 

Figure 4-5. AFM image of the surface of the In-rich CIS film with Na dosing (W590, 

Cu/In = 0.9): (a) 10 μm×10 μm, (b) 1 μm×1 μm. 

The matrix regions of both Cu-rich CIS films are textured with the long axis in [0 1 –1] 

direction of the substrate as was the morphology for the Cu-rich CIS films grown at low 

temperature. The island dimensions, however, are on the order of 3 μm×20 μm and the distance 

between islands is on the order of 20 to 30 μm. In contrast, typical dimensions of the island 

structure for the Cu-rich, epitaxial CIS films grown at lower temperature was 2×8 μm and the 

interspacing between island structures was on the order of 10 μm. Therefore, it is obvious that 

higher growth temperature increased the mobility of adatoms, which resulted in a lower density 

of nucleated islands to produce more elongated island structures. Also, from the similar island 

density and the dimensions of the island structure for the Cu-rich CIS samples with and without 

Na (W558 and W584), it appears that the inclusion of Na did not increase the mobility of 

adatoms in the growth process. 

The increased growth temperature also made it possible to grow nearly island-free 

epitaxial CIS films with an In-rich composition as shown in Figures 4-3 and 4-5 (W560 without 

Na-dosing and W590 with Na-dosing). From these images and SAD patterns presented in 

sections 1.5.2 and 1.5.4, it is concluded that the In-rich CIS films (W560 and W590) grew 

epitaxially. The epitaxial growth of CIS film with an In-rich composition was not observed at 

lower growth temperature of 360 °C. These In-rich CIS films showed an almost island-free 

matrix structure. The island structure shown in Figure 4-3 was observed with a density of less 

than 100 cm-2 and almost no island structure was observed in the In-rich CIS film with Na-dosing 

(W590). Therefore, it is suggested that at high growth temperature, the mobility of adatoms is 

sufficient to allow layer-by-layer growth without the assistance of a Cu-Se secondary phase. It is 

noted that the density of islands at a given growth temperature increases as the Cu/In ratio 

increases. This suggests that the excess Cu-Se is important to the island nucleation and growth 

process. 

59
The faceted island structure of the Cu-rich CIS films (W558, W584) and the island 

shown in the slightly In-rich CIS film without Na dosing (W560) are likely a mixture of a Cu-Se 

secondary phase and CIS as was shown for the Cu-rich CIS films grown at low temperature. Na 

dosing appears to increase the extent of the segregation of Cu-Se secondary phase in the island 

region, which will be confirmed by TEM analysis in section 1.5.3. 

The matrix regions of both In-rich CIS films (W560, W590) are also textured in [0 1 -1] 

direction of the substrate. The roughness of this textured structure in In-rich CIS films at high 

growth temperature, however, is smaller (less than 5 nm) than that (80 nm) in Cu-rich CIS films 

grown at lower temperature. In the case of Cu-rich CIS films at lower growth temperature, the 

formation of an undulated textured surface in the matrix region was attributed to the stress 

relieving process by forming the (1 1 2) CIS planes with the lowest surface energy. For In-rich 

CIS films (W560, W590) grown at high temperature, however, formation of a very thin buffer 

layer at the interface between the film and GaAs is suggested. The existence of this buffer layer 

was proposed from the observation of a small feature in the XRD spectra shown in Figure 4-12 

and the TEM-EDX line scan that shows considerable diffusion of Cu atoms into the GaAs 

substrate (Figure 4-24). The heteroepitaxial stress in the In-rich CIS films (W560, W590) grown 

at high temperature is apparently relieved by this compliant thin buffer layer, which results in the 

shallower undulations on the matrix surface. 

4.3.1 EPMA Analysis 

The compositions of the matrix and the island regions of each film was probed by EPMA 

to determine if there is a difference between the two regions, as observed in lower temperature 

growth. The results of the analysis are summarized in Table 4-2. 

Table 4-2. EPMA analysis for the matrix and island regions of CIS film. 

Sample 

Matrix Region Island Region 

Cu 

(at.%) 

In 

(at. %) 

Se 

(at. %) Cu/In Cu (at. %) 

In 

(at. %) 

Se 

(at. %) Cu/Se 

W558 26.41 23.67 49.92 1.12 58.31 0.33 41.36 1.41* 

W560 25.96 23.78 50.26 1.09 59.25 0.68 40.08 1.48* 

W584 26.09 24.11 49.81 1.08 60.60 0.61 38.79 1.56* 

W590 26.34 23.48 50.18 1.12 59.25 0.68 40.07 1.48* 

*

: Cu/Se ratio after subtracting CuInSe2 composition. 

In this analysis, an accelerating voltage of 6 keV was used. Similar to the lower 

temperature samples, the segregation of Cu-Se into the island region was observed for all CIS 

films, regardless of their bulk Cu to In metal ratio and Na dosing. It also appears that there was 

Cu-Se phase segregation near the surface region of the matrix region. EPMA is more surface 

sensitive (penetration depth of 200 nm at this accelerating voltage), and the Cu to In and Se to In 

ratios were higher than the bulk average determined by ICP. 

60
To determine whether excess Cu-Se exists as a secondary phase or a solid solution and 

whether it resides in the matrix or island regions of the CIS films, SEM micrographs were taken 

after treating each CIS film with a 10% KCN solution to etch any Cu segregates and Cu-Se 

secondary phases exposed on the film surface. The SEM and AFM micrographs shown in 

Figures 4-6 to 4-9 are images of the same samples shown in Figures 4-2 to 4-5, but after etching. 

Note that the images are taken at different magnifications. It is observed that the surface 

morphology of the matrix region of each sample does not change with etching. The grains in the 

island regions of the Cu-rich CIS film without Na dosing (W558), however, appear to be covered 

with a thin layer of a Cu-Se secondary phase. Comparing Figure 4-6b to 4-2b, it is evident that 

etching resulted in the exposure of grains in the island region. 

5 μm 

[0 1 -1] 

2 μm 

(a) (b) 

Figure 4-6. SEM image of the surface of the Cu-rich CIS film (W558, Cu/In = 1.10) after 

etching: (a) 3,000×,(b) island region at 10,000×. 

5 μm 500 nm 

[0 1 -1] 

(a) (b) 

Figure 4-7. SEM image of the surface of the In-rich CIS film (W560, Cu/In = 0.99) after 

etching: (a) 3,000X, (b) island region at 30,000X. 

Interestingly, it is observed that the faceted single grain displayed in Figure 4-4b for the 

Cu-rich, Na-dosed CIS film (W584), was completely removed by etching (Figure 4-8b). Thus a 

Cu-Se phase formed a separate grain as the top part of the island structure of the Cu-rich, Nadosed CIS film (W584). 

61
20 μm 2 μm 

[0 1 -1] 

(a) (b) 

Figure 4-8. SEM image of the surface of the Cu-rich CIS film with Na dosing (W584, 

Cu/In = 1.10) after etching: (a) 1,000X, (b) island region at 10,000X. 

Z = 20 nm/div 

Z = 20 nm/div 

[0 1 -1] 

(a) (b) 

Figure 4-9. AFM images of the surface of the In-rich CIS film with Na dosing (W590, 

Cu/In = 0.9) after KCN etching: (a) 10 μmX10 μm, (b) 1 μmX1 μm. 

For the In-rich CIS film without Na dosing (W560), a large crater was exposed after 

etching at the location of the island structure. Thus the entire island region consists of a Cu-Se 

secondary phase in this case. The EPMA analysis for the island region of Cu-rich, Na-dosed CIS 

film (W584) listed in Table 4-2 suggests that the Cu-Se secondary phase is Cu2-xSe with x ~ 0.5, 

which will be confirmed by TEM analysis (Figure 4-29 in section 1.5.3). The crystal structure, as 

determined by TEM analysis, of the Cu-Se secondary phase in the island region will also be 

reported in section 1.5.3. 

From the AFM images of In-rich, Na-dosed sample (W590) after etching, it was observed 

that the RMS surface roughness was increased to 5.83 nm from the RMS value of 3.25 nm 

before etching. This result indicates that the surface of the In-rich, Na-dosed sample (W590) was 

covered with a thin layer of the Cu-Se phase. 

62
4.3.2 Cross-sectional SEM Images 

To identify the location of the island and the matrix regions and the possible growth 

mode for each CIS film, cross-sectional SEM micrographs were taken on samples prepared by 

cleaving the film/GaAs wafers in the [0 1 -1] direction. 

The SEM cross sectional images in Figures 4-10 and 4-11 clearly indicate that CIS films 

grew uniformly regardless of their Cu/In ratio and the addition of Na at this high growth 

temperature. 

500 nm 500 nm 

GaAs 

CIS 

matrix 

Island 

GaAs 

CIS 

matrix 

(a) (b) 

Figure 4-10. SEM cross-sectional images of CIS films without Na-dosing: (a) Cu-rich 

CIS film (W558, Cu/In = 1.10), (b) In-rich CIS film (W560, Cu/In = 0.99). 

1 μm

1 μm 

Island 

GaAs 

CIS 

matrix 

GaAs 

CIS matrix 

(a) (b) 

Figure 4-11. SEM cross-sectional images of CIS films with Na-dosing: (a) Cu-rich CIS 

film (W584, Cu/In = 1.09), (b) In-rich CIS film (W590, Cu/In = 0.90). 

CIS films, however, showed different growth modes depending on the Cu/In ratio. Curich CIS films (W558 and W584) were grown epitaxially by the S-K growth mode, while the 

nearly stoichiometric or In-rich CIS films (W560 and W590) showed layer-by-layer growth 

mode. From the fact that the layer-by-layer growth mode was observed for both In-rich samples 

(W560 and W590), the mobility of adatoms was apparently sufficient for layer-by-layer growth 

at this growth temperature, and did not require the assistance of the excess Cu-Se or Na-dosing. 

63
4.4 XRD Analysis of CIS Films Grown at High Temperature 

4.4.1 XRD θ−2θ Scan and ω-rocking Curve Analysis of CIS Films Grown at High 

Temperature 

To determine the film orientation and phase, high resolution XRD θ−2θ scans were taken 

for each film with the results summarized in Figure 4-12. The XRD spectra for all CIS films 

showed XRD peaks for only CIS (00n) planes, which is consistent with the epitaxial growth. 

There was also a small feature at 2θ values slightly greater than that for the GaAs (400) 

reflection. To estimate the strain in all CIS films, and to help identify the origin of this feature, a 

XRD space map, which is a series of ω−2θ scans with varying ω values, was performed using 

the X’Pert system. Unfortunately, it was not possible to separate the feature from the intense 

(400) reflection. The ratio of the intensity of this feature to that of GaAs was very small (~ 8×10­

4

) and the signal from the unknown peak can be overwhelmed by the background intensity. More 

importantly, for the crystal plane with higher index usually used in XRD space map, the intensity 

from the feature is very small when there is a strain distribution which is observed from the 

broad shape of the peak as in this case. It is due to the less coherence of the crystal planes with 

high index than (00n) planes when there is a strain in the film. 

Figure 4-12. XRD θ−2θ scan of each CIS film. 

 The XRD θ−2θ scan of In-rich CIS film (W560) after etching by KCN solution shows 

the same feature in the spectrum, which indicates the reflection is not from the Cu-Se secondary 

phase. 

64
Figure 4-13. XRD θ−2θ scan of the In-rich CIS film without Na-dosing (W560) after 

KCN etching. 

As the bump is located at a higher 2θ value in the [001] direction of the reciprocal space, 

the phase or region of the film that is responsible for reflection has a smaller c constant than CIS 

or larger in-plane lattice constant than that of GaAs. Therefore, a possible hypothesis is that a 

very thin GaAs region is near the interface between the film and the substrate is strained relative 

to the film. It is also possible that the broad peak is derived from an InXGa1-XAs layer or CuIn1­

xGaxSe2 layer formed by the inter-diffusion of In or Ga atoms between the film and the substrate. 

The inter-diffusion of In or Ga will be discussed in section 1.3 

To compare the crystal quality of epitaxial CIS films with different Cu/In ratios and Na 

dosing, ω-rocking curve scans for the (008) plane were recorded for each CIS film. As shown in 

Figure 4-14, the FWHM values of the rocking curves for the Cu-rich CIS film (W558 or W584) 

is slightly larger than those for the In-rich CIS films (W560 or W590). This difference in FWHM 

values is likely attributed to the defects at the boundaries between the matrix and the island 

regions in Cu-rich CIS films since these samples have a higher density of islands than the In-rich 

CIS films (W560 and W590), which are nearly island-free. The FWHM value of the substrate 

(35 arcsec) is considerably smaller than any of the CIS films (in range of 648 to 923 arcsec). 

Thus the structural quality of the CIS films is considerably less than the GaAs substrate, 

regardless of the presence of Na or Cu/In ratio. It is concluded that the mobility of the adatoms is 

sufficient at this high temperature and the Cu-Se secondary phase is not needed for the epitaxial 

growth. Similarly, the higher FWHM values for CIS films with Na-dosing (W584, W590) than 

the CIS films without Na-dosing (W558, W560) also indicate that the inclusion of Na did not 

improve the crystal quality and the surface mobility. The relatively large difference in lattice 

constants between GaAs and CIS is likely responsible for the presence of a larger defect density 

and residual strain that dictate the overall structural quality. 

65
4.4.2 Strain Analysis of CIS Films Grown at High Temperature 

An analysis to estimate the change of lattice parameters or the strain of each CIS film was 

performed. A space map analysis was performed on the (228) crystal plane of CIS film and the 

(224) crystal plane of GaAs for each film. The same X’pert system and procedures described in 

chapter 2 were was used for the space map analysis. A hybrid 4-bounce monochromator with 

four Ge (220) single crystals and 1/2 inch slit were used to produce the incident beam, and a 

triple axis, including two of the Ge (220) single crystals, was used to collect the diffracted beam. 

The combined resolution for this ω−2θ scan is 0.0001°. The values of the (ω, 2θ) coordinates, 

were obtained by manually maximizing the intensity of the (228) peak of CIS film and that of the 

(224) peak of GaAs, respectively. 

Figure 4-14. ω-rocking curve of the CIS (008) plane for each film and the GaAs (400) 

plane. 

In the XRD space map measurement, only the Cu Kα1 emission line with a wavelength 

of 1.54089 Å was used. From this wavelength and the lattice parameter of GaAs (5.654 Å from 

JCPDS32-0389), the theoretical value for the (ω, 2θ) coordinate set of the GaAs (224) plane is 

calculated as ω = 6.6146° and 2θ = 83.758°. After considering the offset between the theoretical 

(6.6146, 83.758) and the measured (6.6466, 83.822) values for the GaAs (224) plane, the values 

for the CIS (228) crystal plane were calculated and summarized in Table 4-3. Also, from the 

estimated values presented in Table 4-3, the lattice constants of each film were calculated and 

listed in Table 4-4. These values were compared to the unstrained lattice parameters of CIS film 

(a = 5.782 Å, c= 11.619 Å from JCPDS40-1487) to give the lattice constants of each film 

relative to the GaAs substrate. 

66
Table 4-3. Estimated values of ω, 2θ, φ, and psi for the GaAs (224) and CIS (228) 

planes. 

ID GaAs CIS 

2θ ω φ psi 2θ ω φ psi 

W558 83.7580 6.6146 -0.18 0.79 81.1875 5.2932 -0.52 0.78 

W560 83.7580 6.6146 -0.22 0.96 81.1947 5.3075 -0.48 0.45 

W584 83.7580 6.6146 0.81 0.22 81.1977 5.2990 0.81 0.61 

W590 83.7580 6.6146 -0.45 0.27 81.2017 5.2475 -0.45 0.56 

Table 4-4. Lattice constants and lattice mismatch relative to GaAs for each CIS film. 

ID k008 (Å-1) k220 (Å-1) c(Å) a(Å) mismatch 

W558 0.6893 0.4880 11.606 5.796 2.51% 

W560 0.6894 0.4880 11.604 5.796 2.51% 

W584 0.6894 0.4881 11.604 5.791 2.42% 

W590 0.6889 0.4888 11.613 5.786 2.33% 

A plot of the lattice constants of each CIS film in reciprocal space clearly shows that each 

CIS film grown at this high temperature was nearly fully relaxed (Figure 4-15). In this case, the 

lattice mismatch between each CIS film and the GaAs substrate might have been accommodated 

by misfit dislocations in these films, especially near the boundaries of the island structures 

embedded in the matrix region for the Cu-rich CIS films (W558, W584) and near the interface 

between the CIS film and GaAs substrate for the In-rich CIS films (W560, W590). As suggested 

previously in this section, the XRD spectra shown in Figures 4-12 and 4-13 show a small and 

broad peak positioned at a slightly greater 2θ value than the reference (400) GaAs reflection. It is 

proposed that a thin buffer layer was formed in the GaAs near the interface between GaAs and 

CIS film. This buffer layer is thought to be made from the diffusion of Cu, In or Ga atoms to 

form a graded buffer layer near the interface. This buffer layer is assumed to present an in-plane 

lattice constant close to that of CIS to permit epitaxial growth of a nearly relaxed film. 

67
Figure 4-15. Lattice constants of each CIS film in the reciprocal space. 

4.5 TEM Analysis of CIS Films Grown at High Temperature 

As observed in the SEM micrographs in section 1.3, Cu-rich CIS films (W558, W584) 

showed the island structures embedded in the smooth matrix region. Also, the top part of the 

island structure in Cu-rich CIS film with Na-dosing (W584) was suggested as Cu1.5Se from the 

result by EPMA analysis in section 1.3.1. Therefore, it is necessary to investigate the crystal 

structure of each CIS film more extensively at high spatial resolution. 

For this purpose, TEM analysis was performed in selected areas on each CIS film. 

Selected area diffraction (SAD) is used to measure the local lattice spacing. For the assignment 

of each diffraction spot observed using SAD patterns, the distance between the transmitted beam 

and each diffraction spot is first measured. Then, the inter-planar spacing for each diffraction 

spot is calculated by the following equation: 

Rdhkl = λL , (4-1) 

where R, dhkl, λ, and L are the measured distance between the diffraction spot and the transmitted 

beam, the inter-planar spacing of the corresponding diffraction spot, wavelength of the electron 

beam, and the camera length of the microscope (82 cm) used to obtain the diffraction pattern, 

respectively. The wavelength of the electron beam at the accelerating voltage of 200 kV used in

this study is 0.0252 Å. The values of dhkl, is then calculated from Equation 4-1 and the 

corresponding crystal plane (h k l) is determined by comparing the inter-planar spacing value, 

dhkl with the theoretical or experimental values. 

68
In TEM analysis, dark field images are often taken to yield an image with better contrast 

and to better demark the dislocations in the film. Practically, a dark field image is obtained by 

locating the diffracted beam at the optic axis using a two-beam condition, where diffraction is 

obscured from only two crystal planes. Therefore, the dark field image under this condition 

shows the feature in the film that has orientation parallel to the diffracted beam, and those 

features that have an orientation (h k l) normal to the diffracted beam are not shown in the dark 

field image. This approach can be used to determine the orientation of a dislocation in the 

material or to identify an area with a different orientation or crystal structure from other areas. 

4.5.1 TEM Analysis of the Cu-rich, Na-free CIS Film (W558) 

Though the epitaxial growth of all CIS films were confirmed by XRD analysis as in 

section 1.4, it is necessary to identify the orientation of the CIS film relative to that of GaAs. 

SAD patterns of film and substrate for the Cu-rich film without Na addition were taken using 

selected area aperture with a diameter of 500 nm at [1 -1 0] zone axis as in Figure 4-16. 

B 

(11-1) 

(111) 

(220) 

(002) 

B 

(11-2) 

(112) 

(220) 

(004) 

(a) (b) 

Figure 4-16. SAD pattern of GaAs substrate and CIS region of Cu-rich CIS film without 

Na (W558) at zone axis [1 -1 0]: (a) GaAs, (b) CIS. 

The lattice constants of CIS film were calculated from the pattern shown in Figure 4-16b 

and a value of 82 cm for the camera length. The value of λL is calculated from the product of the 

values of the separation distance and the electron beam energy. This is only possible, however, 

when the sample is positioned exactly at the eucentric position. As in the strain analysis using 

XRD data, to estimate the offset in the λL value, it is assumed that the in-plane lattice constant of 

GaAs is the fully relaxed value since it is considerably thick (~ 400 μm) than the overgrown CIS 

film (~ 0.5 μm). The distance between the transmitted beam (indexed with B in Figure 4-16) and 

the spot for the (220) plane in the SAD pattern for GaAs area (Figure 4-16a) is 10.7 mm. 

Therefore, to produce the in-plane lattice constant of GaAs (a = 5.654 Å), the required value of 

λL is 21.39 mm·Å. For the CIS area, the distance of the (220) spot from the transmitted beam is 

10.45 mm, and results in the value of d(220) = 2.047 Å (21.39 mm·Å/10.45 mm) and the in-plane 

lattice constant, a = (8)1/2d(220) = 5.790 Å. The values for the inter-planar spacing, dhkl, and their 

crystal plane assignments are summarized in Table 4-5. 

69
Table 4-5. Lattice parameters of GaAs and CIS calculated from SAD pattern at [1 -1 0] 

zone axis for Cu-rich CIS film without Na (W558). 

Distance to the spot (mm) Inter-planar spacing, d (Å) Lattice parameter (Å) 

(220) (111) or (112) d(220) d(111) or d(112) a c 

GaAs 10.70 6.56 1.999 3.261 5.654 5.637 

CIS 10.45 6.35 2.047 3.369 5.790 11.86 

* Estimated 0.1 mm uncertainty. 

The assignments to specific crystal planes are labeled in Figure 4-16. From the lattice 

constants presented in Table 4-5, this CIS film was grown in a nearly relaxed state. A SAD 

pattern was also taken at the region near the interface region between CIS and GaAs to check if 

the Cu-rich CIS film (W558) shows pseudomorphic growth. A selected area (SA) aperture with a 

diameter of ~ 250 nm was used, where half of the SAD aperture was filled with CIS material and 

the other half was filled with GaAs material. As the pattern in Figure 4-17 shows, double spots 

are apparent for each crystal plane, and then the CIS film did not grow pseudomorphically near 

the interface region. 

These results are consistent with the result of the XRD strain analysis presented in the 

previous section. However, from the observation of the diffused spots in Figure 4-17, it is 

assumed that the SAD spots in Figure 4-17 might be a mixture of the spots from the bulk region 

of CIS and GaAs materials and the spots from the buffer layer region since the size of the SAD 

aperture is too big for a thin buffer layer. Therefore, the growth model is postulated as follows. 

At the initial stage of the growth, a thin buffer layer with an in-plane lattice close to that of CIS 

film is formed in the GaAs region near the interface. With the help of this buffer layer and 

resulting decrease of the heteroepitaxial strain, the CIS film grows epitaxially by 2-D layer 

growth mode. However, there is also the segregation of Cu-Se phase from the Cu-rich 

composition at some of the nucleation sites, which results in the faster growth and thicker island 

region. 

The spots labeled as 1 & 1` in Figure 4-17, correspond to the family of (2 2 0) crystal 

planes, {2 2 0}, and those labeled as 2 & 2`are for the {1 1 1} family. 

Dark field image were taken by using (0 0 -4) diffraction vector and the images of the 

island and matrix regions are presented in Figure 4-18. The images in Figure 4-18 show that the 

island region consists of an underlying matrix layer and a multiple grain upper region. The image 

shows the underlying matrix region contains several dislocations originating at the CIS/GaAs 

interface. The image in Figure 4-18b similarly shows the matrix region contains dislocations. 

The dislocations extend from the substrate to both in the island and matrix regions. 

Therefore, it is proposed that the heteroepitaxial strain due to the lattice mismatch between the 

CIS film and GaAs was largely relieved by the formation of the misfit dislocations during the 

initial stage of the growth. 

70
B 

2 & 2` 

1 & 1` 

Figure 4-17. SAD pattern at the interface of the Cu-rich CIS film without Na addition 

(W558) and the GaAs substrate at zone axis [-1 1 0]. 

500 nm 500 nm 

CIS 

island region 

CIS 

matrix 

region 

GaAs 

(a) (b) 

Figure 4-18. Dark field images of the Cu-rich, no Na sample (W558) using g = (0 0 -4) 

diffraction vector at 50,000X magnification: (a) island region, (b) matrix region. 

SAD patterns from a single grain and from multiple grains in the island region were taken 

(Figure 4-19). These patterns suggest that, as in the case of Cu-rich CIS film grown at lower 

temperature (See section 2.5 for sample W367), twins are formed, where each grain was grown 

as a single crystal with a small rotation or mis-orientation relative to each other. The reduction of 

the heteroepitaxial strain is also attributed to the existence of this twin structure. 

71
B 

B 

(a) (b) 

Figure 4-19. SAD pattern for a single grain and multiple grains in the island region at 

zone axis [1 -1 0]: (a) Single grain, (b) Multiple grains. 

4.5.2 TEM Analysis of In-rich, Na-free CIS Film (W560) 

The same TEM analysis was performed on the In-rich, Na-free CIS film (W560), to 

determine the values of lattice constants and the relative orientation of CIS region with respect to 

the substrate. Figure 4-20 shows SAD patterns of CIS and GaAs regions of the sample at zone 

axis [-1 1 0]. As shown in this figure, the SAD patterns for both GaAs and CIS film show the 

same symmetry and orientation at the zone axis [-1 1 0] but different distances between 

corresponding different spots. 

The same calibration procedure in section 1.3.1 was performed that requires a (GaAs) =

5.654 Å for the 82 cm camera length. It was calibrated for the (2 2 0) spot and the λL value was 

determined as 20.65 mm·Å. By using this value, the lattice constants of CIS region and GaAs 

substrate were again calculated and the values are summarized in Table 4-6. In addition, each 

diffraction spot was assigned to the specific crystal plane as labeled in Figure 4-20. 

(0 0 2) 

(1 1 1) 

(2 2 0) 

B 

(0 0 4) 

(1 1 2) 

(2 2 0) 

B 

(a) (b) 

Figure 4-20. SAD patterns of the GaAs substrate and CIS region for In-rich, Na-free CIS 

film (W560) at zone axis [-1 1 0]: (a) GaAs, (b) CIS. 

72
Table 4-6. Lattice parameters of GaAs and CIS region for In-rich CIS film (W560) 

calculated from SADP at [-1 1 0] zone axis. 

Distance to the spot (mm) Inter-planar spacing, d (Å) Lattice parameter (Å) 

(220) (002) or (004) d(220) d(002) or d(004) a c 

GaAs 10.330 7.300 1.999 2.829 5.654 5.658 

CIS 10.167 7.250 2.031 2.848 5.745 11.39 

* Estimated 0.1 mm uncertainty. 

The results shown in Table 4-6 indicate that the in-plane lattice parameter “a” of the Inrich CIS film (W560) is only slightly less than the unstrained material (lattice mismatch = 

1.61%). The island-free, In-rich, Na-free CIS film grew heteroepitaxially in a nearly unstrained 

state with good crystal quality. SAD pattern of the interface region between GaAs and CIS at 

zone axis [-1 1 0] is shown in Figure 4-21. 

2` 

2 

1` 

1 

B 

3 

3` 

Figure 4-21. SADP near the interface between CIS film and GaAs substrate, W560. 

The diffraction spots were indexed by comparing with the theoretical values and are 

labeled in Figure 4-21. The diffraction spots labeled 1, 2 and 3 in the figure are produced from 

the CIS region and 1`, 2` and 3` are from the GaAs region. Their assignments are summarized in 

Table 4-7. 

Table 4-7. Assignment of diffraction spots for the GaAs and CIS 

regions near the interface. 

1 2 3 

GaAs {1 1 1} family {0 0 2} family {2 2 0} family 

CIS {1 1 2} family {0 0 4} family {2 2 0} family 

* Estimated 0.1 mm uncertainty. 

Values of the lattice parameters were calculated from the measured distance between 

each spot and the transmitted beam (camera length = 137 cm, λL = 52.77 mm·Å The lattice 

constants are summarized in Table 4-8. The in-plane lattice constant, “a”, of the CIS near the 

73 
interface is almost the same as unstrained CIS, which indicates that the film grew in the nearly 

relaxed state from the interface. 

Table 4-8. Lattice parameters of GaAs and CIS near the interface calculated from SAD 

patterns at [-1 1 0] zone axis. 

Distance to the spot(mm) Inter-planar spacing, d (Å) Lattice parameter (Å) 

(220) (002) or (004) d(220) d(002) or d(004) a c 

GaAs 26.4 18.70 1.999 2.822 5.654 5.644 

CIS 26.0 18.2 2.030 2.899 5.742 11.60 

* Estimated 0.1 mm uncertainty. 

Figure 4-22 shows a bright field and a dark field image of W560 using (0 0 -4) diffracted 

beam. Dislocations are observed in the dark field image of the CIS region, some of which extend 

to the interface between CIS and GaAs, and release the strain energy that would be produced 

from the lattice mismatch between CIS and GaAs. 

The possible inter-diffusion of Cu, In or Ga was proposed in section 1.4 from the 

observation of the low intensity, broad peak at slightly higher value of 2θ than the GaAs (400) 

peak (Figure 4-12). To probe this possibility, TEM-EDX scan was performed across the 

interface. A bright field image of the scan region and TEM-EDX scan are shown as in Figures 4­

23 and 4-24, respectively. 

500 nm 500 nm 

Pt 

coating 

Pt 

coating 

CIS CIS 

GaAs GaAs 

(a) (b) 

Figure 4-22. Bright field and dark field images of the In-rich, Na-free sample (W560) 

using g = (0 0 -4): (a) bright field image, (b) dark field image. 

74
500 nm 

GaAs 

CIS 

TEM EDX-scan 

Figure 4-23. Bright field image of the TEM-EDX scan region near the GaAs/CIS 

interface of In-rich, Na-free sample (W560). 

From the EDX scan shown in Figure 4-24, a considerable amount of Cu diffusion from 

the CIS film into the GaAs was observed. In Figure 4-24, the position of the interface was 

determined where Se signal sharply decreases to zero. It is known that Cu atom diffuses rapidly 

in GaAs even at low temperature by an interstitial diffusion process [4-2]. The diffusion 

coefficient of Cu in GaAs was reported to be as D = 3×10-2 cm2

s

-1× exp(-0.53/kT) and 7.641×10­

6

 cm2/s at 464 °C [4-3]. The solubility was found also to be high (5.66×1014cm-3 at 464 °C) [4-3]. 

Therefore, it is proposed that the bump on the right side of GaAs (400) peak in the XRD spectra 

is due to Cu diffusion from CIS film to the GaAs during the growth, and the increase of the inplane lattice constant is attributed to the formation of Cui defects in the GaAs substrate. 

Additionally, the broad feature was not observed in the XRD spectra of CIS films grown at low 

temperature. However, there was no diffusion of other elements such as Ga or In atoms 

observed. 

75
Figure 4-24. TEM-EDX scan of In-rich, Na-free CIS film (W560) near the interface. 

4.5.3 TEM analysis of the Cu-rich, Na-dosed CIS film (W584) 

It is recalled from the results shown in section 1.2 that the top part of the faceted island 

structure in the Cu-rich, Na-dosed CIS film (W584) was completely removed after KCN etching 

(Figure 4-8). The conclusion that the upper layer in the island facets are primarily a Cu-Se 

secondary phase was further supported by the EPMA results (Table 4-2). The island region of the 

sample was probed using TEM to independently investigate the upper and lower parts of the 

island region. 

Figure 4-25 shows the bright field images of the island and matrix regions of this CIS film 

(W584). The EPMA analysis (Table 4-2) suggested that the top part of the island region has 

composition close to Cu1.5Se. Figure 4-26 shows the SAD patterns for both GaAs (a) and the 

island bottom area (b) that is CIS in the Cu-rich CIS film with Na-dosing (W584). 

76
200 nm 

GaAs 

CIS matrix 

1 μm 

GaAs 

CIS 

Matrix 

Island 

Bottom 

CIS 

Island Top 

Cu1.5Se 

(a) (b) 

Figure 4-25. Bright field images of the Cu-rich, Na-dosed CIS film (W584): (a) island 

structure, (b) matrix region. 

B 

(311) 

(020) 

B 

(312) 

(020) 

(a) (b) 

Figure 4-26. SAD pattern of GaAs and matrix regions of the In-rich, Na-dosed sample 

(W584) at [-1 0 3]: (a) GaAs, (b) matrix region. 

As observed in Figure 4-26, the CIS film was again grown epitaxially with the same 

orientation as the underlying GaAs substrate. The lattice constants of the CIS film were 

calculated and summarized in Table 4-9. The results indicate that Cu-rich CIS, Na-dosed CIS 

film (W584) was again grown nearly relaxed epitaxial, regardless of Na-dosing. 

Table 4-9. Lattice parameters of GaAs and CIS calculated from SADP at [-1 0 3] zone 

axis. 

Distance to the spot(mm) Inter-planar spacing, d (Å) Lattice parameter (Å) 

(020) (311) or (312) d(020) d(311) or d(312) a c 

GaAs 7.375 12.25 2.827 1.702 5.654 5.556 

CIS 7.240 18.20 2.880 1.733 5.760 11.26 

* Estimated 0.1 mm uncertainty. 

77 
The double spots in the SAD pattern of the interface region are displayed in Figure 4-27., 

the CIS film appears to have grown in a nearly relaxed state in Figure 4-27. This is the same 

finding for the film grown without Na-dosing (W558 and W560). 

B 

(311) 

(020) 

D 

D 

D 

D 

Figure 4-27 .SAD pattern of the interface between GaAs and CIS film at [-1 0 3]. 

The double spots denoted as D correspond to the spots from both GaAs and CIS film. 

Figure 4-28 shows the SAD patterns taken for the upper and lower areas in the island region. The 

lattice constants of each phase, presumably Cu1.5Se for the upper layer and CIS for the lower 

layer, were calculated from the spot distances measured on patterns shown in Figure 4-28. From 

the results shown in Table 4-10, the in-plane lattice constant of the Cu1.5Se phase was strained 

relative to the CIS phase below it. 

B 

(311) (220) 

B 

(312) (220) 

(a) (b) 

Figure 4-28. SAD pattern of island region of the Cu-rich, Na-dosed sample (W584) at [-1 1 2] 

zone axis: (a) upper layer of island region (Cu1.5Se), (b) lower area of the island region (CIS). 

78
Table 4-10. Lattice parameters of GaAs and CIS calculated from SADP at [-1 0 3] zone axis 

rL = 20.65. 

Distance to the spot 

(mm) 

Inter-planar spacing, d 

(Å) 

Lattice parameter (Å) 

(220) (311) or (312) d(220) d(311) or d(312) a c 

Upper layer 

(Cu1.5Se) 

10.10 6.150 2.045 3.358 5.784 5.882 

Island bottom 10.10 6.225 2.045 3.317 5.784 11.34 

(CIS) 

* Estimated 0.1 mm uncertainty. 

To verify whether the Cu-Se phase in the top part of the island structure is a Cu1.5Se 

phase, TEM-EDX line scan was performed along the depth direction of the island structure as in 

Figure 4-29. As shown in Figure 4-29, it is observed that the Cu/Se ratio of the Cu-Se region in 

the top part of the island structure where there is no In, is close to 1.5. 

Figure 4-29. TEM-EDX line scans of W584 along the depth of the island structure. 

The coexistence of Cu1.5Se phase with CIS in the epitaxial growth of Cu-rich CIS (Cu/In 

= 1.4 ~ 2.6) was reported before by Fons et al. [4-4]. In their work, they reported the existence of 

the wedge-shaped epitaxial Cu1.5Se phase on the surface of the epitaxial CIS grown on (001) 

GaAs. They calculated the lattice constant of the Cu1.5Se phase using XRD space map and 

obtained an in-plane lattice constant of 5.770 Å and an out-of plane lattice constant of 5.722 Å. 

They also grew epitaxial Cu1.5Se on (001) GaAs substrate without CIS phase and reported the inplane lattice constant of 5.764 Å as a nearly completely relaxed lattice constant. The in-plane 

79
lattice constant calculated here for the Cu1.5Se phase found in the Cu-rich CIS film with Nadosing (W584) is, therefore, close to the value reported in their study. Also, as in their study, the 

in-plane lattice constant of Cu1.5Se was strained to that of CIS. 

From the fact that the segregation of the epitaxial Cu1.5Se was not observed in the Cu-rich 

CIS film without Na-dosing (W558) with the same Cu/In ratio in section 1.5.1, it is proposed that 

Na dosing resulted in the segregation and the epitaxial growth of the Cu1.5Se phase on top of the 

epitaxial CIS phase in the island structure by acting as a surfactant during the growth. Also, Nadosing was known to improve the crystal quality by suppressing the formation of the defect pairs 

in the film. Therefore, it is postulated that this suppression of defects in CIS film by Na might 

suppress the existence of Cu-Se nano-crystalline phase in the CIS region and results in the 

segregation of the Cu-Se phase in the form of Cu1.5Se. 

4.5.4 TEM Analysis of In-rich, Na-dosed CIS film (W590) 

The final sample examined by TEM in the set was grown at the high temperature of 464 

°C with In-rich composition and Na-addition. The SEM images and XRD analysis showed this 

sample grew island-free and nearly relaxed. The dark field images shown in Figure 4-30 were 

taken using the g = (0 2 2) diffracted beam. 

CIS 

Matrix 

CIS 

Island 

GaAs 

500 nm 

300 nm 

CIS matrix 

GaAs 

(a) (b)

Figure 4-30. Dark field images of In-rich, Na-dosed CIS film (W590): g = (0 2 2): 

(a) Dark field image: Island (50,000X), (b) Dark field image: Matrix (100,000X). 

There are a few grain boundaries observed in the island region, which may be the result 

of the formation of twins. Also, the voids and dislocations are observed at the distinctive 

interface between the matrix region and the GaAs substrate. The dislocations or grain boundaries 

from the possible twin structures in the island region are believed to release the heteroepitaxial 

strain in this film. The SAD patterns (Figure 4-31) were taken for the matrix and substrate region 

and the spots were assigned to crystal planes as in Figure 4-31. 

80
B (111) 

(200) 

B 

(112) 

(200) 

(a) (b) 

Figure 4-31. SAD pattern of GaAs and matrix regions of In-rich, Na-dosed CIS film 

(W590) at [0 1 -1] zone axis: (a) GaAs, (b) matrix region (CIS). 

Again, the matrix region of the CIS film showed the same orientation as that of GaAs and 

the epitaxial growth. The lattice parameters of GaAs and the CIS region were also calculated as 

in Table 4-11. 

Table 4-11. Lattice constants of GaAs and CIS calculated from SADP at [0 1 -1] zone. 

Distance to the spot(mm) Inter-planar spacing, d (Å) Lattice parameter (Å) 

(200) (111) or (112) d(200) d(111) or d(112) a c 

GaAs 7.48 6.48 2.828 3.264 5.656 5.648 

CIS 

matrix 

7.30 6.33 2.897 3.341 5.794 11.54 

* Estimated 0.1 mm uncertainty. 

The lattice parameter values of the CIS film listed in Table 4-11 show that In-rich, Nadosed CIS film also grew epitaxially and almost completely relaxed similar to the other CIS 

films grown at high growth temperature. The double spots in Figure 4-32 indicate that In-rich, 

Na-dosed CIS film grew epitaxially and with almost no strain. 

81
B 

D 

D 

(111) 

(200) 

D 

D 

Figure 4-32. SAD pattern of In-rich, Na-dosed CIS film (W590) at the interface between 

GaAs and CIS film. 

Diffraction spots for twins were observed in the SAD pattern of the island structure 

(Figure 4-33), and it is suggested that the formation of these twins is induced by the 

heteroepitaxial stress (lattice mismatch = 2.26%). The formation of twins in the island region 

leads to the epitaxial growth in this film by relieving the heteroepitaxial stress. 

B 

(112) 

(200) 

T T 

Figure 4-33. SAD pattern of island region above the matrix region of In-rich, Na-dosed 

CIS film (W590): ‘T’ means diffraction patterns from twins. 

82
4.6 Conclusions 

The Cu-rich CIS films (W558 and W584) showed morphology similar to those of Cu-rich 

CIS films grown at the low temperature. This morphology consists of a 2-D matrix layer with the 

nucleation of isolated island regions. The thickness of the matrix under the island region is 

greater than the island-free area. In contrast, a two-dimensional growth mode resulting a very 

smooth surface was observed for In-rich CIS films (W560 and W590), regardless of Na dosing. 

This is also in contrast to the polycrystalline structure observed for the In-rich CIS film grown at 

low temperature. It is suggested that the surface mobility of adatoms was increased for 2-D 

epitaxial growth at a high growth temperature. The assistance of Cu-Se secondary phase appears 

not necessary for the epitaxial growth at this high growth temperature. 

The In-rich CIS films (W560, W590) showed very few island structure (100 cm-2), while 

Cu-rich CIS films (W558, W584) showed highly faceted island structures, which increased their 

size slightly with Na dosing (from 10 ~ 15 μm to 15 ~ 20 μm). However, the spacing between 

island structures was nearly the same for Cu-rich CIS films regardless of Na-dosing. XRD θ−2θ

scans indicated that all CIS films were grown heteroepitaxially. XRD space maps also indicate 

that all CIS films grew in a nearly fully relaxed state, which also occurred at the lower growth 

temperature (Chapter 2). 

From the TEM-EDX line scan for the cross-section of In-rich CIS film without Nadosing (W560), it is observed that there is a considerable Cu-diffusion from the film to the 

substrate, which is possibly due to the higher growth temperature and possibly attributed to the 

sequential deposition process of the PMEE reactor allowing longer migration step for the 

adatoms. It is proposed that this Cu diffusion formed a thin buffer layer with a larger in-plane 

lattice parameter in the GaAs substrate near the interface region. It is also suggested that the 

feature of slightly higher value of 2θ than the GaAs (400) peak in the XRD spectra is due to the 

decrease of c-axis lattice constant (higher 2θ value or lower d004) as well as the increase of inplane lattice constant in that layer. It is postulated that this buffer layer with increased in-plane 

lattice constant partly relieved the heteroepitaxial stress between CIS film and the GaAs 

substrate and enabled CIS films to grow epitaxially without strain along with the strain relieving 

process by dislocations. 

The difference of the growth process between Cu-rich and In-rich CIS films is postulated 

as follows. As the liquidus of Cu2-xSe was suggested as 525 °C in Fon’s et al.’s work [4-4], it is 

possible that the Cu-Se secondary phase in the Cu-rich CIS films at the higher growth 

temperature of 464 °C exist as a liquid-like state during the growth. As the surface mobility is 

enough at this temperature without the assistance of the Cu-Se secondary phase, the adatoms 

form a 2-D layer from the initial stage of the growth with a buffer layer and resulting smaller 

lattice mismatch at the interface. For In-rich CIS films, part of the heteroepitaxial strain that is 

not relieved by the formation of the buffer layer is suggested to be relieved by the formation of 

dislocations at the interface as the film continues to grow with the same growth rate for the entire 

2-D layer region. However, for Cu-rich CIS films, the Cu-Se phase is segregated at some of sites 

where the dislocation is concentrated, and then results in the faster growth rate by even lowering 

the surface energy with a surfactant effect at the sites, and the formation of the thicker island 

regions over the sites. 

83
The formation of the epitaxial Cu1.5Se phase at the top part of island structure was 

observed for the Cu-rich CIS film with Na dosing (W584). It is postulated that the Na results in 

the segregation of the epitaxial Cu1.5Se phase in the top part of the island region by suppressing 

the formation of the defects, and the existence of the Cu-Se phase as a nano-crystalline phase in 

the CIS film. This segregation was not observed for the Cu-rich CIS film without Na-dosing 

(W558) with the same Cu/In ratio. 

The SAD patterns taken at the interface of CIS film and GaAs substrate showed double 

spots for all CIS films. It confirms that all CIS films grew with relaxed state with in-plane lattice 

constant close to the original value. Twins were also observed in the island region of the Cu-rich 

CIS film without Na-dosing (W558) and the island region of the In-rich CIS film with Na-dosing 

(W590) by TEM-SADP. It is proposed that the heteroepitaxial stress due to the lattice mismatch 

leads to the formation of those twins in the island region. The composition of the Cu-Se phase 

observed in the island region of Cu-rich CIS film with Na-dosing (W584) was confirmed as 

Cu1.5Se by TEM-EDX line scan of the island structure. It is also identified as epitaxial to the CIS 

film in the island structure. The in-plane lattice constant of this Cu1.5Se phase was calculated

from the TEM-SADP analysis and its value, 5.784 Å was close to that of CIS (5.784 Å). 

4.7 References cited 

[4-1] D.W. Niles, K. Ramanathan, F. Hasoon, R. Noufi, B.J. Tielsch, and J.E. Fulghum, J. Vac. 

Sci. Technol. A 15, 3044 (1997). 

[4-2] F.C. Frank, and D. Turnball, Phys. Rev. 104, 617 (1956). 

[4-3] R.N. Hall, and J.H. Racette, J. Appl. Phys. 35, 379 (1964). 

[4-4] P. Fons, S. Niki, A. Yamada, and H. Oyanagi, J. Appl. Phys. 84, 6926 (1998). 

84
5 Study of the Defect Structure of CGS by 

Photoluminescence Spectroscopy 

5.1 Introduction 

 CuGaSe2 has recently received attention as a material for the top cell in a tandem 

structure solar cell since it has high band gap (1.73 eV at 10 K). Therefore, its defect structure 

determines the material’s electrical properties. From the previous studies, it is known that CGS 

films are always grown p-type since the compensating donor state, GaCu, has a much deeper 

energy level (130 meV) than that of the corresponding donor in CIS (25 meV). The acceptor 

state, VCu, in CGS has a similar energy level (50 ~ 60 meV) as that of CIS (40 ~ 50 meV). Also, 

shifting to Ga-rich compositions increases the Cu-deficiency and the resulting increase in the 

VCu, which is an acceptor. Thus the CGS top cell does not possess a natural buried junction as is 

thought to exist in CIGS cells. Being able to type convert CGS would be a major contribution. 

Therefore, it is useful to study its defect structure with varying Cu/Ga ratio and without influence 

from grain boundaries. 

In this study, the defect structure of CGS films was investigated using the same CGS 

films studied in the previous chapter (Cu-rich and Ga-rich composition). In addition the effect of 

Na of adding Na during growth was examined to determine its impact on the defect structure of 

CGS, particularly on the VCu or CuGa defect states. Cu and Ga-rich CGS films with and without 

Na-dosing were also characterized by photoluminescence (PL). Films were also analyzed by PL 

after annealing in Se to investigate the effect of Se concentration on the defect structure as well 

as the transition related to the VSe. 

5.2 Theoretical Background 

Photoluminescence (PL) is a process of photon emission by the transition between two 

defect states in the forbidden gap or between one of the defect states and band edge (conduction 

band (CB) or valence band (VB)) after the photon energy is absorbed by the material. When light 

with photon energy higher than the band gap of the material is absorbed, electron-hole pairs are 

generated in the material. These electron-hole pairs are usually bound as pairs by electrostatic 

forces, and are called excitons. When these excitons move freely without being bound to a state 

in the band gap, they are called free excitons. An exciton does not contribute to the current and it 

can have a different energy level as it is excited and finally breaks into free electron and free 

hole. The peaks corresponding to these excitonic transitions are usually observed near the band 

edge in a PL spectrum, the exact position depending on their excitation states. The binding 

energy of the exciton can be calculated using the following expression for each excitation energy 

level n [5-1] 

∗ 4 

E = m 

2 

r q 

2 2 , (n=1, 2, 3...), (5-1) ex 2h ε n 

where mr

*

 is the reduced mass defined as mr

*

 = me

*

mh

*

/(me

*

 + mh

*

) with me

*

 and mh

*

 being the 

85
electron and hole effective masses, respectively. ε is the dielectric constant, h is Planck’s 

constant, and n is an integer and equal to 1 for the exciton ground state. The binding energy of 

the exciton can be calculated from literature data for me

*

 and mh

*

 or alternatively, it can be 

calculated by using the following equation if the peak energies for the ground state (EFX1) and 

the first excited state (EFX2) of the free exciton were known. 

4 EFX = (EFX 1 − EFX 2 ) (5-2) 3 

Sometimes, a shoulder is observed at the lower energy side of the excitonic emission 

peak. If the intensity of this shoulder becomes smaller, the shoulder is usually attributed to the 

bound exciton emission peak, and can be separated from the excitonic emission peak by 

measuring the PL emission with varying excitation power. In the excitation power dependent PL 

spectrum, both peaks show a super linear dependence of PL intensity on the excitation power, 

which is proof of excitonic recombination [5-2]. The emission intensity from excitonic 

recombination is proportional to the product of the number of electrons in CB and holes in the 

VB. As the photons generate electron-hole pairs, the emission intensity shows a quadratic 

dependence on the excitation power in principle. The number of bound excitons in thermal 

equilibrium with the free ones is also proportional to the square of the number of incident 

photons, or excitation power. As the free excitons are in thermal equilibrium with bound 

excitons, the temperature dependence of the integrated excitonic emission intensity from both 

excitons can be represented as follows, 

NFE = C1e−Ea1 / kT N BE , NFCP = C2

' 

e−Ea 2 / kT NFE = C2 e−( Ea1 +Ea 2 ) / kT NBE , (5-3) 

where NBE(T), NFE(T), and NFCP(T) are the numbers of bound excitons, free-excitons and free 

carriers, respectively, k is Boltzman constant, Ea1 is the localization energy of excitons to a 

bound center and Ea2 is the dissociation energy of a free exciton to a free electron-hole pair. As 

the total number of particles is in equilibrium, the total number of bound and free excitons and 

carriers is proportional to the excitation power, Pexc 

NBE + NFE + NFCP = CPexc , (5-4) 

where C is a constant. If the emissions from both bound and free excitons are observed, the 

integrated PL intensity, I(T), can be represented by the following equation: 

I (T ) = r1 NBE + r2 NFE 

CPexc NBE + NFE + N FCP , (5-5) 

where r1 and r2 are the ratios of the collected emission through the aperture to the total emission 

by bound excitons or free excitons, respectively. Finally, the integrated intensity can be 

represented as: 

1+ rC e−Ea1 / kT 

I (T ) = C0

1+ C1e−Ea1 / kT +

1 

C2e −( Ea1 +Ea 2 ) / kT (5-6) 

86
with C0 = r1CPexc, C2 = C1×C2

`

, and r = r2/r1. Therefore, from parameter fitting of the 

experimental data, the dissociation energy of a free exciton can be obtained and compared to 

known values. The localization energy of the bound exciton can also be obtained, from which the 

activation energy of the bound center, EBC can be calculated by using Hayne’s rule [5-3], Ea2 = 

CEBC with C ~ 0.1. 

There are usually other peaks present in a PL spectrum, which are often related to the 

donor-acceptor (D-A) pair transition. They usually exist within 200 meV from the band edge for 

CIS or CGS. The peak energy of this transition is given by the following expression [5-4]. 

2 

E = Eg − (E + Ed ) +

e 

+ f (r) (5-7) a 4πε r ε 0 r 

In this equation, Eg is a band gap, Ea is the acceptor state energy level, Ed is the donor 

state energy level, εr is the dielectric constant of the material, ε0 is the dielectric constant of the 

vacuum, r is the separation of the donor-acceptor pair involved in the transition. The last term, 

f(r) is a quantum mechanical correction, which can be neglected if r is large compared to the 

radii of the acceptor and donor states. If the radii of those states overlap, the position of the 

emission energy increases by quantum mechanical repulsion, which can be calculated from 

solution of the stationary Schrödinger equation for those states. The D-A related transition shows 

a peak shift to higher energy (blue shift) with increasing excitation power as the average distance 

between the D-A pair decreases with excitation power. 

The ionization energies of the states involved in a D-A pair transition can be estimated by 

the temperature dependent PL measurement [5-2, 5-5] as follows. First, the efficiency of the 

donor-acceptor recombination for the donor state is represented by 

j 

ηj = PDA , (5-8) DA P j + P j + P j + P j 

Dh DA DO TD 

where each term represents recombination with valence band holes with probability of 

occurrence Pj

Dh, the recombination with holes at the acceptor states with probability PjDA, 

recombination by other processes with probability Pj

DO, and thermal ionization to the conduction 

band with probability Pj

TD, respectively. Also, the efficiency of donor-acceptor recombination 

for the acceptor state, is given by 

j AD η AD = j j

P j

j j (5-9)

PeA + PAD + PAO + PTA 

with Pj

eA representing the transition probability with electrons at the conduction band, PjAD the 

probability of donor-acceptor recombination with electrons at donor state, Pj

AO the probability of 

other processes and Pj

TA the probability of thermal ionization to the valence band edge. The 

overall efficiency of the donor-acceptor recombination is given by the product of equations (5-8) 

and (5-9), 

87
j j 

ηj ηj = PDA PAD (5-10) DA AD P j + P j + P j + P j P j + P j + P j + P j 

Dh DA DO TD eA AD AO TA 

The integrated intensity of the emission related to the donor-acceptor recombination is 

given by, 

I = j η A , (5-11) DA 

where A is a constant. As the thermal transition probabilities Pj

TA and PjTD are given by an 

exponential function of ionization energy, they can be represented as, 

PTA

j = BA exp(−EA / kT ) , (5-12) 

where EA and ED are the ionization energies of the acceptor and donor states, respectively. As 

only Pj

TA and PjTD are functions of temperature and all recombination probabilities are 

independent of temperature, equation (5-11) can be written as follows: 

C3 I = [1+ C4 exp(−ED / kT )][1+ C5 exp(−EA / kT )] , (5-13) 

where all recombination probabilities are shown as constants. Therefore, the ionization energies 

of the donor and acceptor states can be estimated by parameter fitting of the measured integrated 

emission intensity of the related peak using equation (5-13). 

5.3 Experimental Approach 

Samples prepared in the MEE reactor were characterized by PL using the facilities at 

Yungnam University at Kyongsan, Korea. The system is a commercial unit (SPEX750M). The 

PL measurements were carried out using the 514 nm line of an Ar-ion laser with a beam 

diameter of ~100 μm and an excitation power of 2 mW. PL spectra were first obtained at 10 K to 

probe possible excitonic emissions near the band edge as well as any donor-acceptor transitions. 

The sample temperature was varied from 10 to 300 K and the temperature-dependent PL 

intensities measured to allow calculation of the activation energies of defect states involved in a 

particular transition. In addition to characterizing the defect structure of films grown in the MEE 

reactor, the effect of subsequent annealing in a Se overpressure was performed to better 

understand the atomic origin of the defect. These PL measurements were performed on samples 

annealed by placing a solid Se source in the chamber with the as-grown films at 300 °C for 30 

min. PL measurements were not performed for CIS films since a photomultiplier tube was the 

only available detector and thus suitable for the higher band gap energy films. 

The CGS samples studied were grown by MEE. The samples were selected to study the 

effect of Na-dosing on the point defects in CGS. Two CGS films were grown with Cu-rich 

(Cu/Ga = 1.19, W552) and Ga-rich (Cu/Ga = 0.96, W567) compositions. For the study of the Na 

effect on the defect structure, Na was incorporated into each film during growth using a Knudsen 

cell delivering NaF. 

88
5.4 PL Measurements of Epitaxial CGS Films 

5.4.1 PL Spectra 

PL measurements were performed on the as-grown CGS films and CGS films after Se 

annealing at 10 K and the resulting spectra are shown in Figures 5-1 through 5-8. All spectra 

show a distinctive peak near 1.5 eV (between 1.498 and 1.508 eV). This peak is assigned to the 

transition of excitons bound to copper centers with trigonal symmetry in Cu-doped GaAs by 

Wang et al. [5-6] and they suggested that the CuGa – VAs defect pair is responsible for this center. 

Therefore, this observation confirms the diffusion of Cu atoms into the GaAs substrate in all 

CGS films, and supports the hypothesis proposed in Chapter 4 that a diffusion layer forms within 

the GaAs substrate. 

Figure 5-1. PL spectrum of Cu-rich, Na-free CGS film (W552) at 10 K. 

89
Figure 5-2. PL spectrum of Ga-rich, Na-free CGS film (W567) at 10 K. 

Figure 5-3. PL spectrum of Cu-rich, Na-doped CGS film (W583) at 10 K. 

90
Figure 5-4. PL spectrum of Ga-rich, Na-doped CGS film (W599) at 10 K. 

Figure 5-5. PL spectrum at 10 K of Cu-rich, Na-free CGS film (S552) after Se annealing. 

91
Figure 5-6. PL spectrum at 10 K of Ga-rich, Na-free CGS film (S567) after Se annealing. 

Figure 5-7. PL spectrum at 10 K of Cu-rich, Na-dosed CGS film (S583) after Se annealing. 

92
Figure 5-8. PL spectrum at 10 K of Ga-rich, Na-doped CGS (S599) after Se annealing. 

The peak position in the range 1.673 and 1.679 eV in each PL spectrum, except in the 

Ga-rich, Na-free CGS film (W567) spectrum, is assigned to the conduction band to acceptor state 

transition based on previous studies [5-7, 5-8, 5-9]. From this previous work as well as the peak 

energy difference (51 ~ 57 meV) from the band edge (1.730 eV at 10 K), the acceptor state 

involved in this transition is assigned to VCu acceptor state. It is also observed that the peak at 

1.646 or 1.647 eV in the PL spectra of the Cu-rich CGS films (both Na-dosed and Na-free) 

disappeared after Se annealing (compare Figures 5-1, 5-3, 5-5, and 5-7), which strongly suggests 

the assignment of this peak to a VSe-related transition. 

The peak located at 1.667 eV in the PL spectrum of Ga-rich, Na-free CGS film, W567 

(Figure 5-2) also disappeared after Se-annealing. Therefore, it is also assigned to a VSe-related 

transition. The peak positioned at 1.690 eV in the PL spectrum of the Ga-rich, Na-free CGS film 

(W567) appears to be related to an excitonic transition. It is not assigned to this specific 

transition since there is no transition with the corresponding energy difference. For the Ga-rich, 

Na-dosed CGS film (W599), the PL spectrum (Figure 5-4) included a peak at 1.593 eV, which is 

assigned to the VCu-VSe D-A transition as previously suggested in the literature [5-5, 5-8]. The 

fact that the ratio of this peak to the peak at 1.676 eV assigned to VCu – CB transition decreases 

after Se-annealing for all CGS films, is further evidence that this peak involves a VSe – related 

transition. 

The peak for the VCu-CB transition (at 1.676 eV) that was found in the Ga-rich Na-free 

CGS film (W567) and the peak at 1.690 eV were not observed in the Ga-rich, Na-dosed CGS 

film. The disappearances of the VCu-CB transition in the Ga-rich, Na-dosed CGS film (W599) is 

likely due to the decrease of VCu from the formation of NaCu with Na addition. Also, the 

disappearance of the exciton related peak at 1.690 eV is likely due to the poorer crystalline 

93
quality of the Ga-rich, Na-dosed CGS film (sample W599 has the highest FWHM in the HRXRD 

patterns with the lowest Cu/Ga ratio of 0.81) than that of Ga-rich CGS film. Also, a new peak 

appears at energy of 1.421 eV in the PL spectrum when Na is added to the Ga-rich CGS film 

(Figure 5-8). Gislason et al. proposed that the PL peak at 1.41 eV observed in GaAs doped with 

Cu and Li atoms is related to the CuGa-Lii-related transition [5-10]. 

Figure 5-10 is a TEM-EDX line scan of the Cu-rich, Na-dosed CGS film (W583) without 

Se-annealing for the region shown in the micrograph in Figure 5-9. Considerable diffusion of Na 

into the GaAs substrate is evident in the profile shown in Figure 5-10. These profiles are 

consistent with the assignment of observed peak intensities to Na and Cu related defect in GaAs. 

It is also noted that considerable diffusion of Na and Cu into the substrate occurs when CGS is 

replaced by CIS for an In-rich composition and of course Na-dosing (Figures 5-11 and 5-12). 

400 nm 

CGS 

GaAs 

Scan direction 

Figure 5-9. TEM-EDX line scan region of the Cu-rich, Na-dosed CGS film (W583). 

94
Figure 5-10. TEM-EDX line scan of Cu-rich, Na-doped CGS film (W583) before Se 

annealing. 

500 nm 

Scan 

direction 

CIS 

GaAs 

Figure 5-11. TEM-EDX line scan region of In-rich, Na-dosed CIS film (W590). 

95
GaAs 

CIS 

Figure 5-12. TEM-EDX line scan of In-rich, Na dosed CIS film (W590) before Se 

annealing. 

Therefore, it is suggested that this peak at 1.41 eV is related to a CuGa-Nai-related 

transition in GaAs, which forms upon both Cu and Na diffusion during growth into the GaAs 

substrate from the film. 

There is a new peak observed at 1.656 eV in the PL spectrum of the Cu-rich, Na-dosed 

CGS film after Se-annealing (S583) as indicated in Figure 5-7. This peak is believed to be 

related to the transition involving a Na-related defect. Assuming that Vcu is present in the 

material, even for Cu-rich conditions, the search for a transition of a defect state pair with a 

relatively small formation energy and defect level would include VGa-VSe: 1.646 eV, VCu-CB: 

1.678 eV, VCu-VSe: 1.6 eV, Cu(Ga)–VSe: 1.667 eV. Since these defect pairs are accounted for, a 

Na-related defect is the likely possibility for this peak 

5.4.2 Thermal Quenching of PL Peaks 

To investigate the PL peaks involving only donor-acceptor transitions, a selected 

transition was thermally quenched to probe the activation energy of the defect state of the 

shallower energy level according to the following equation [5-11]: 

I = C1 (5-15) (1+ C2 exp(−E / kBT )) 

96
Figure 5-13. Thermal quenching of PL peak at 1.647 eV in the Cu-rich, Na-free CGS film 

(W552). 

Figure 5-14. Thermal quenching of PL peak at 1.646 eV from Cu-rich, Na-dosed CGS 

film (W583). 

97
Figure 5-15. Thermal quenching of PL peak at 1.667 eV from Ga-rich, Na-free CGS film 

(W567). 

Figure 5-16. Thermal quenching of PL peak at 1.656 eV from Cu-rich, Na-dosed CGS 

film after Se annealing (S583). 

In equation (5-15), I is the integrated intensity, C1 and C2 are constants, kB is the 

Boltzman constant, and E is the activation energy of the shallower defect state of the two. 

The results of the thermal quenching studies are graphically shown in Figures 5-13 to 5­

16. Several conclusions can be drawn from these results. First, analysis of the thermal quenching 

data for the peak at 1.647 eV found in the Cu-rich, Na-free CGS films (W552, W583) shown in 

Figures 5-13 and 5-14 produces slopes of (29±2.1) and (21±1.4) meV, respectively. These 

energy levels are likely related to the VGa defect state, since they are close to the value (20 meV) 

98
reported in a previous study [5-12]. Therefore, the peak at 1.647 eV is assigned to the VGa-VSe 

transition. An activation energy of (10±0.34) meV was derived from the thermal quenching data 

of the PL peak at 1.667 eV (see Figure 5-15). From the fact that this peak disappeared after Se 

annealing and the value of the activation energy (10 meV) reported for the CuGa defect by 

Bauknecht et al. [5-13], this transition is assigned to the CuGa-VSe transition. 

The spectrum shown in Figure 5-7 reveals a new peak at 1.656 eV for the Cu-rich, Nadosed CGS film after Se-annealing (S583) The activation energy of this transition was calculated 

as (20±0.78) meV (Figure 5-14), and it was tentatively assigned to the NaGa acceptor state. It is 

also suggested that this peak is related to the NaGa–VSe transition, and this conclusion is based on 

the following rationale. First, the activation energies estimated for the VGa defect from the data 

shown Figures 5-13 and 5-14, are 29 and 21 meV. In addition, the reported activation energy of 

the defect level of VSe is 80 meV [5-8]. From these energy levels, the Coulombic term in 

equation (5-7) for the peak at 1.647 eV is estimated as 18 and 26 meV, respectively. This value 

shows good agreement with Coulombic energy of 26 meV in the report of Bauknecht et al. [5­

13] with an estimated typical distance between the two charged defects of 4 nm. Using these 

values to approximate the Coulombic energy of the defect complex responsible for the peak at 

1.656 eV, the donor state involved in this transition is suggested to be a VSe since the activation 

energy for the donor state is calculated as 72 ~ 80 meV. The peak energies observed for each 

sample in this study and their assignments are summarized in Tables 5-1 and 5-2. 

Table 5-1. Peak energies and their assignments observed in CGS films in this study. 

Peak Energy (eV) Cu-rich Ga-rich Cu-rich, Na-dosed Ga-rich, Na-dosed 

1.427 

1.500 CuGa – VAs CuGa – VAs CuGa – VAs CuGa – VAs 

1.580 VCu – VSe + 

1.600 VCu – VSe VCu – VSe VCu – VSe 

1.647 VGa – VSe VGa – VSe 

1.656 

1.667 CuGa – VSe 

1.676 VCu – CB VCu – CB 

Table 5-2. Peak energies and their assignments observed in CGS films after Se 

annealing in this study. 

Peak Energy (eV) Cu-rich Ga-rich Cu-rich, Na-dosed Ga-rich, Na-dosed 

1.427 CuGa – Nai 

1.500 CuGa – VAs CuGa – VAs CuGa – VAs 

1.580 VCu – VSe + 

1.600 VCu – VSe VCu – VSe VCu – VSe 

1.647 

1.656 NaGa – VSe 

1.667 

1.676 VCu – CB VCu – CB VCu – CB 

99
5.5 Conclusions 

The point defects of CGS films grown by MEE on GaAs were characterized by low 

temperature PL spectroscopy. In particular, the effects of the Cu/Ga ratio, Na dosing, and 

subsequent Se-annealing on the defect states were investigated. Peak assignments to specific 

defects were suggested by comparing the spectra to previously reported results, observing 

changes in the spectra as the film composition was changed, and estimating activation energies 

estimated from the results of thermal quenching experiments. This work supports the findings of 

previous studies of the defects VCu, VGa, CuGa, and VSe with corresponding activation energies of 

50 ~ 60, 20 ~ 30, 10, and 75 ~ 80 meV. In addition, a new defect state was observed at 20 meV 

above the valence band edge from the thermal quenching of the peak at 1.656 eV in the Cu-rich, 

Na-dosed CGS film after Se annealing. For the Ga-rich, Na-dosed CGS film, a new PL peak was 

observed at 1.427 eV in the PL spectrum after Se annealing and assigned to the transition from 

the exciton bound to the CuGa – Nai defect pair. 

Composition profiles across the CGS/GaAs interface (TEM_EDX) showed substantial 

diffusion of Cu into the GaAs as well as Na when added to the film. A peak near 1.5 eV was 

observed in all CGS films. Consistent with previous studies on GaAs, this peak was assigned to 

the transition from the excitons bound to the CuGa – VAs defect pair, consistent with diffusion of 

Cu into the GaAs substrate in all CGS films. The diffusion of Na and Cu into the GaAs substrate 

was observed by TEM-EDX line scan of In-rich, Na-dosed CIS film (W590). From this 

observation and the report of the CuGa-Lii-related transition peak in the GaAs doped with Na and 

Cu, it is proposed that the peak at 1.421 eV observed in the PL spectrum of Ga-rich, Na-dosed 

CGS film after Se annealing (S599) is related to the CuGa-Nai-related transition. 

A peak at 1.67 eV was assigned to the VCu-CB transition by comparison of the energy 

difference from the band edge with the values reported in the previous studies. Activation energy 

in the range 21 to 29 meV was estimated for the peak at 1.647 eV from the thermal quenching 

experiments, and it was assigned to the activation energy of the VGa defect, again from 

comparison with the value in previous studies. From the disappearance of the peak at 1.647 eV 

after Se-annealing, this peak was assigned to the VGa-VSe transition. 

An activation energy of 10 meV was calculated from the thermal quenching studies as the 

energy level for a defect state involved in the transition responsible for the 1.667 eV peak in the 

PL spectrum of Ga-rich, Na-free CGS film. This energy level was assigned to the CuGa defect 

state from the comparison with the results of previous studies and the peak was assigned to the 

CuGa-VSe transition since it also disappeared after Se-annealing. 

Another new defect state at 20 meV above the valence band was observed from the 

thermal quenching of the peak at 1.656 eV in the PL spectrum of the Cu-rich, Na-dosed CGS 

film after Se-annealing (S583). This energy level was tentatively assigned to the NaGa acceptor 

state, and this peak was assigned to the VSe-NaGa transition from considerations of the activation 

energy and the energy difference from the band edge including estimation of Coulombic energy. 

100
From the PL measurement results for the Cu-rich CGS films, it is concluded that the VCu 

is the dominant defect state, even in the Cu-rich CGS films since it has a low activation energy 

(50 ~ 60 meV) and the lowest formation energy (0.66 eV, first principles calculation in Table 5­

3) of all defect states in CGS. It is also postulated that the disappearance of the peak from the 

VGa – VSe transition in the Cu-rich CGS films after Se annealing is a result of a decreased 

stability of VGa defect state due to the reduction of VSe concentration for the formation of the 

stable defect pair. 

 The GaCu defect state was not observed in all CGS films as a result of its deep energy 

level (130 meV) and high formation energy (3.22 eV, Table 5-4). Therefore, it is expected that 

Na-dosing may not increase the hole concentration of the CGS films to the extent that it does in 

CIS films since the concentration of the compensating donor, GaCu in CGS films is much lower 

than that of InCu in CIS films. This is a result of a large difference in their activation energies (25 

meV for CIS and 130 meV for CGS). 

Formation of a n-type buried junction near the surface region might be achieved by the 

formation of a very thin layer of Cu(In,Ga)Se2 with high In/Ga ratio, low Cu/(In+Ga) ratio, and 

with sufficient Se flux near the surface region should lead to the n-type layer from the formation 

of InCu shallow donor states in the layer with less VGa concentration. Of course, this will lower 

the band gap energy of this layer and modify the absorption process. 

5.6 References cited 

[5-1] J.I. Pankove, Optical Processes in Semiconductors (Dover, New York, 1971), p. 12. 

[5-2] A. Yamada, Y. Makita, S. Niki, A. Obara, P. Fons and H. Shibata, J. Appl. Phys. 79, 4318 

(1996). 

[5-3] J.R. Haynes, Phys. Rev. Lett. 4, 361 (1960). 

[5-4] H.B. Bebb and E.W. Williams, Semiconductors and Semimetals in Transport and Optical 

Phenomena (Academic Press, New York, 1972) p. 182. 

[5-5] A. Yamada, P. Fons, S. Niki, H. Shibata, A. Obara, Y. Makita and H. Oyanagi, J. Appl. 

Phys. 81(6), 2794 (1997). 

[5-6] Z.G. Wang, H.P. Gislason and B. Monemar, J. Appl. Phys. 58, 230 (1985). 

[5-7] S. Chichibu, Y. Harada, M. Uchida, T. Wakiyama and S. Matsumoto, S. Shirakata, and S. 

Isomura, H. Higuchi, J. Appl. Phys. 76, 3009 (1994). 

[5-8] J.H. Schön, F.P. Baumgartner, E. Arushanov, H. Riazi-Nejad, C. Kloc, and E. Bucher, J. 

Appl. Phys. 79(9), 6961 (1996). 

[5-9] J.H. Schön, and E. Bucher, Solar Energy Materials & Solar Cells 57, 229 (1999). 

[5-10] H.P. Gislason, Z.G. Wang, and B. Monemar, J. Appl. Phys., 58 (1), 240 (1985). 

101
[5-11] P.J. Dean, Phys. Rev., 157, 655 (1967). 

[5-12] J.H. Schön, J. Phys. D : Appl. Phys. 33, 286 (2000). 

[5-13] A. Bauknecht, S. Siebentritt, A. Gerhard, W. Harneit, S. Brehme, J. Albert, S. Rushworth, 

and M.Ch. Lux-Steiner, Thin Solid Films 361-362, 426 (2000). 

102
6 Investigation of Deep-Level Defects in Cu(In,Ga)Se2 

Solar Cells by Deep-Level Transient Spectroscopy 

6.1 Introduction 

It is well known that CuInGaSe2 (CIGS) has a large homogeneity range, and such 

deviation from stoichiometry in compound semiconductors is often attributed to antisite defects, 

vacancies, interstitials, and defect clusters in the material. It is important to understand these 

properties to further improve the cell efficiency and reduce the cost of solar energy. Deep level 

defects play an important role in determining the recombination and trapping mechanisms (and 

hence the minority carrier lifetimes) in a semiconductor. Deep level transient spectroscopy 

(DLTS) is a high frequency transient capacitance technique and is the most sensitive technique 

for defect detection. Zhang [6-1,6-2] has calculated the transition energy for a large number of 

defects in CIS. However, there are some discrepancies between his calculation and experimental 

observations. According to the PL measurement by Schoon [6-3], InCu should be a shallower 

donor than Cui and has transition energy about 0.2 eV below conduction band. In Zhang’s 

calculation, the anion site defect is neglected. However, many other researches believe that VSe 

is the shallowest donor level [6-3] with energy level of transition Ec- 0.06 eV and results in a 

significant lattice relaxation during the ionization process. Knowledge of the defect energies as 

well as their densities are important inputs to developing meaningful device models. DLTS is a 

standard diagnostic technique for determining the trap properties such as trap energy level, 

capture cross section, and trap concentration. 

6.2 DLTS Principles and Experimental Procedures 

Deep level defects play an important role in determining the recombination and trapping 

mechanisms (and hence the minority carrier lifetimes) in a semiconductor. The deep level 

transient spectroscopy (DLTS) is a high frequency transient capacitance technique and is the 

most sensitive technique for defect detection [6-4]. For a n+

-p (CdS-CIGS) junction diode, the 

space charge region extends mainly into the p-region, and the local charges are due to negatively 

charged ionized acceptors. If a forward bias is applied, the electrons will be injected into region. 

Once the minority electrons are trapped in a defect level, the net negative charges in the SCR 

will increase. This in turn will reduce the width of SCR, and causes a positive capacitance 

change. Thus, the DLTS signal will have a positive peak. If holes are injected into the SCR and 

captured by the majority carrier traps, the DLTS signal will have a negative sign. Figure 6-1 

shows the deep level filling sequence. 

A boxcar integrator [6-5] uses a fast sampling technique to obtain an average of the 

repeated signal transients on some set time which we will call t1. The double boxcar has the 

capacity to sample the signal at times, which are termed t1 and t2. An analog circuit is used to 

integrate or average the value of ΔC over the two fixed times which are called windows. Figure 

6-2 shows how the rate window works in the analysis of DLTS data. The left side of the figure 

6-2 indicates the response of the transient signal, it can be written as 

103
Δ ( ) = ΔC(0)exp( − n (6-1) Ct et ) 

If we define time constant τ = 1/en, the above equation can be written as 

ΔC t( ) = ΔC(0)exp( −t / ) τ (6-2) 

The right side is the resulting DLTS output signal for this fixed rate window. It is 

ΔC t() − Δ () = S τ = ΔC(0)[e −t2 /τ −t

1 /τ Ct () − e ] (6-3) 2 1 

τ To have the maximum of S(τ), we need to have dS() = 0 , and this occurs at dτ

t t

τ = 1 − 2 (6-4) ln( / ) t t 1 2 

Thus, the emission rate corresponding to the maximum of a trap peak observed in a DLTS 

thermal scan is a precisely defined quantity. If we have a series of temperature measurements 

using different t1 and t2, one can calculate the activation energy using an Arrhenius plot. 

DLTS is used to measure the transient capacitance change of a junction device after deeplevel traps in the space charge region have been filled with either the majority or minority carrier 

charges, and the trap density and energy levels of these deep level traps can be extracted from 

DLTS measurements. By scanning the transient capacitance over a range of temperature, the 

emission rate as a function of temperature can be obtained. The activation energy for an 

electron trap and capture cross section can be estimated from an Arrhenius plot: 

1 E 

en = =σ nvth Nc exp(− a ) (6-5)

τ kT 

where en is the electron emission rate, σn is the capture cross section, Nc is the electron effective 

density of states and νth is the electron thermal velocity. The electron trap density, NT, can be 

calculated by 

2ΔC NT = Na (6-6) C0 

Equation (6-6) is valid for NT < Na. ΔC is the height of capacitance change due to the 

emission of charge carriers from the trap level, which is proportional to the probed trap 

concentration. Carrier concentration, Na, is determined by the C-V measurements, which give 

the average carrier concentration in the CIGS absorber layer. C0 is the zero bias capacitance that 

is obtained from the C-T data at the corresponding DLTS peak temperature. 

104
6.3 Results and Discussions 

DLTS studies were conducted on three types of thin-film photovoltaic cells: 

(i) CIS cells (denoted “UF CIS cells” in the suite for easy reference) developed by 

the University of Florida using a bilayer precursor process. 

(ii) CIGS cells (denoted “EPV CIGS cells”) developed by Energy Photovoltaics Inc. 

using a physical-vapor deposition process. 

(iii) CIGS cells (denoted “NREL CIS cell”) developed by the National Renewable 

Energy Laboratory using a three-stage precursor process. 

Details of the analysis results are discussed in detail below. 

6.3.1 UF Samples from Binary Bilayer Rapid Thermal Process 

The I-V curve for the UF 5% cell showed a Voc= 0.296 V and fill factor = 49.54%. The 

Jsc for this cell is 34.65 mA/cm2, which is a very good value for CIS cells. The low fill factor is 

largely due to the high sheet resistance. This can be improved by adjusting the Cu:In:Se 

composition ratio, especially Cu to In ratio and and by creating a composition gradient to achieve 

the bandgap grading with the incorporation of Ga. The Voc is largely affected by the mid gap 

defects in the space charge region. To detect the defect level in this 5% cell, a Deep Level 

Transient Spectroscopy (DLTS) measurement was performed on this 5% cell. 

Figure 6-3 shows the C-V measurements on the 5% UF cell at different temperatures. 

The carrier concentration is on order of 1015 cm-3. The C-V data was measured on the same cell 

as that used in the DLTS measurement. 

Figure 6-4 shows the measured capacitance versus temperature for the 5% UF cell, which 

yields the value for C0 at the corresponding temperature of the DLTS peak. The DLTS spectra 

measured for this 5% sample with a reverse bias voltage, VR =-0.5 V, a trap-filling pulse of 0.4 

V, and a saturation pulse width of 10 ms are shown in Figure 4-5. The sample is maintained at 

reverse bias. The 5% sample shows a deep majority carrier trap (negative signal) at a 

temperature around 290 K. The activation energy calculated from the Arrhenius plot as shown in 

Figure 6-6 is 0.54 eV above valence band, which may be attributed to the Cu on In interstitial 

(EA6) according to the diagram. The average carrier concentration at 290 K from the C-V 

measurements is 2×1015 cm-3. Combing the C-T and C-V scans, NT is determined to be 4.6×1012 

-3 cm . 

To detect possible minority traps, the sample was cooled down to 77 K to freeze out all 

the minority traps. The sample was then heated with a pulse amplitude equal to 0.7 V at the 

same reverse set-off voltage of -0.5 V and a saturation pulse width of 10 ms to observe the 

capacitance change due to the minority carrier emission. As expected, a minority carrier trap 

peak was observed at a lower temperature ~200 K as shown in Figure 6-7. This minority carrier 

trap is in addition to the majority trap that was observed. The activation energy is 0.52 eV below 

the conduction band as shown in Figure 6-8. A possible physical origin of this electron trap is 

105
the actually the same as the majority trap, which is Cu on In interstitial [6-1, 2]. The trap density 

is estimated to be 1.3×1012 cm-3. 

Optical DLTS was also performed on this 5% sample. In optical DLTS, a laser with 

wavelength of 532 nm is applied with a reverse bias of 0.5 V. Only minority traps were 

observed in the optical DLTS spectra as shown in Figure 6-9. The activation energy calculated 

for E1 is 0.16 eV shown in Figure 6-10. This trap is attributed to In on Cu site defect InCu. The 

activation energy for the E2 defect at 300 K could not be estimated since the peak shift did not 

follow the trend; the shorter the rate window, the higher the temperature at which the peak 

appears. Likely, this is because the capture cross-section is a thermally activated process, and it 

is related to temperature by [6-4] 

σ σ= 0e 

−ΔE 

kT (6-7) 

6.3.2 EPV CIGS Sample 

The EPV samples are produced by the co-deposition in a physical vapor deposition 

(PVD) system. The DLTS spectra measured from the EPV sample with a reverse bias voltage, 

VR =-0.1 V, a trap-filling pulse of 0.3 V, and a saturation pulse width of 10 ms is shown in 

Figure 6-11. The EPV sample showed a deep majority carrier trap (negative signal) at a 

temperature around 270 K. The DLTS peak shifts to higher temperatures with shorter rate 

windows. Since the CIGS film is p-type material, the majority carrier trap is referred to as a hole 

trap, and the minority carrier trap as an electron trap. The activation energy calculated from the 

Arrhenius plot is 0.94 eV as shown in Figure 6-12, which may be attributed to the Se vacancy 

VSe. Figure 6-13 shows the C-V data measured on the same EPV cell as that used in the DLTS 

measurement. The average carrier concentration at 270 K from the C-V measurements is 3×1015 

cm-3. From the C-T scan (Figure 6-14), C0 in Eqn. (6-6) was found to be 3.01 pf, and NT was 

determined to be 6.5×1013 cm-3. 

6.3.3 NREL CIGS Sample from Three Stage Process 

The absorber from NREL was deposited by three stage process [6-6] in a PVD system. A 

solar cell 19.2% total area efficiency based on Cu(In,Ga)Se2 thin film material has been produced 

via their three stage process. In the first stage, the precursor (In,Ga)2Se3 layer is formed by the 

coevaporation of In, Ga and Se elemental source at substrate temperature about 260 ̊C. In the 

second stage, the Cu and Se are coevaporated to bring the overall composition to near Cu-rich 

regime (Cu/(In+Ga) ~ 0.97 – 1.08) at 560 ̊C. Between the second stage and third stage, the 

substrate is under the Se flux while ramping up the temperature. During the third stage, In, Ga 

and Se are added with the amount one ninth of the first stage to bring the overall composition 

back to Cu-poor at the same substrate temperature as the second stage. Then, the film is cooled 

down to 350  ̊C by Se flux for 20 min. This process requires very precise control of the source 

fluxes. 

A pulse amplitude equal to 0.4 V was applied to the NREL CIGS sample at a reverse setoff voltage of -0.5 V. A minority carrier trap peak was observed at a lower temperature of around 

106
100 K as shown in Figure 6-15. The activation energy is 0.067 eV below the conduction band as 

shown in Figure 4-16. A possible physical origin of this electron trap is the Se vacancy (VSe) [6­

3]. The trap density is estimated to be 4.2×1013 cm-3. The DLTS scan at higher temperature 

showed another possible positive peak, however, the activation energy for this trap could not be 

resolved by heating the sample above 300 K. 

Therefore, it is noted that the minority traps are likely the dominant traps in the NREL 

sample. To detect other possible minority traps, a small forward bias (0.1 V) was applied and the 

sample cooled to 77 K to freeze the minority carriers. The sample was then heated with VR = -0.5 

V, trap-filling pulse amplitude of 0.7 V, and a saturation pulse width of 10 ms to observe the 

capacitance change due to the minority carrier emission. Figure 6-17 shows three minority 

carrier traps with the addition of a small peak around 250 K, which was not apparent in the 

DLTS spectra of the 0.4 V pulse amplitude. The Arrhenius plot in Figure 6-18 showed that the 

activation energy for the trap at 80 K is 0.14 eV. The trap density is estimated to be 2.54×1013 

cm-3. The electron trap observed at 80 K is possibly related to a vacancy on a Se site, which 

might be the same trap level as the one observed with the 0.4 V pulse amplitude as shown in 

Figure 6-15. The activation energy at 335 K could not be estimated since the peak shift does not 

follow the trend; the shorter the rate window, the higher the temperature at which the peak 

appears. 

6.4 Conclusions 

The DLTS and C-V measurements have been performed on UF CIS sample and compared with 

the DLTS measurements on EPV and NREL CIGS samples. The defect activation energy and 

density were determined from these measurements. Table 6-1 summarizes the DLTS data on 5% 

cell by RTP on bilayer precursor. In the UF CIS sample, a middle gap majority trap is detected 

and three other minority traps are also detected. The middle majority trap seems to be the 

dominant defect in the cell. This confirms the observation of the low Voc in I-V curve. 

Therefore, to improve the cell efficiency, the junction quality must be improved by reducing the 

middle gap defect. Adding Ga or S to create the bandgap gradient and adjusting the Cu:In:Se 

composition ratio will be the reasonable methods. The EPV sample showed a majority carrier 

(hole) trap with activation energy and trap density determined. The NREL CIGS sample showed 

three minority carrier traps. Possible defect origins were suggested from this DLTS study. The 

results for samples from both locations are listed in Table 6-2. From the analysis in Tables 6-1 

and 2, it is concluded that VSe and InCu are two main donor levels, and VCu is the shallowest 

acceptor level, and CuIn is the main middle level recombination center which are consistent with 

the information from literature [6-1~3, 6-7]. The free energy associated with the formation of 

some defect structures is so small that little increase in thermodynamic potential results and 

hence there is insufficient driving force to ensure their elimination under many synthesis 

conditions. The formation energy of these defects in CIS according to several authors 

calculations are in the order of VCu<InCu<CuIn ≈ InCu < VSe [6-1~2, 6-7]. However, the defect 

distribution depends significantly on the stoichiometry. 

107
6.5 References cited 

[6-1] S. B. Zhang, Su-Huai Wei, and Alex Zunger, Physical Review B, Vol. 57, No. 16 

pp9642 –9656, 1998. 

[6-2] S. B. Zhang, Su-Huai Wei, and Alex Zunger, Physical Review Letters, Vol. 78, 

No. 21 pp4059 – 4062, 1997. 

[6-3] J.H. Schon and E. Bucher, Solar Energy Materials & Solar Cells, Vol. 57, pp. 

229-237, 1999. 

[6-4] S.S. Li, Semiconductor Physical Electronics, Plenum Press, New York and 

London, 1993. 

[6-5] D.V. Liang, J. Appl. Phys, Vol. 45, pp.3014, 1974. 

[6-6] B. M. Keyes, P. Dippo, W. Metzger, J. AbuShama, and R. Noufi, the 29th IEEE 

Photovoltaic Specialist Conference, New Orleans, LA, 2002. 

[6-7] G. Kühn, W. Möller and W. Engel, Thermochimica Acta, Vol. 78, No. 1-3, pp. 

129-134, 1984. 

108
(a) 

 CuInxGa1-xSe2 

CdS 

V 

t = 0 

0 V = -5 Volts 

-5 

(b) 

 CuInxGa1-xSe2 

CdS 

V 

t = 10 ms 

0 V = 0 Volts 

0 

 CuInxGa1-xSe2 

CdS 

V 

0 

(c) 

t = 10 ms + 

V = -5 Volts 

- 5 

Figure 6-1. DLTS majority trap filling sequence for CdS/CIGS junction (a) 

under reverse bias without pulse (b) pulse injection (c) after pulse injection. 

109
Figure 6-2. Example of how the double boxcar averager is used to define the 

rate window [6-5]. 

110
30 

35 

40 

45 

50 

55 

60 

Capacitance C (pF)

300 K 

270 K 

200 K 

120 K 

-0.2 -0.1 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 

Reversed Bias (V) 

Figure 6-3. Capacitance–Voltage measurement for UF CIS cell. 

111
Figure 6-4. Capacitance-temperature measurement of the 5% cell. 

112
Figure 6-5. . DLTS scan for UF 5% cell at a reverse bias of –0.5 V, a pulse 

height of 0.4 V, and pulse width of 10 ms. 

113
Figure 6-6. The Arrhenius plot of emission rate en/T2

 vs. 1/T. 

114
Figure 6-7. DLTS scan for UF 5% cell at a reverse bias of –0.5 V, a pulse 

height of 0.7 V, and pulse width of 10 ms. 

115
Figure 6-8. The Arrhenius plot of emission rate en/T2

 vs. 1/T for DLTS scan at 

reverse bias 0.5V, pulse height 0.7 V and pulse width 10 ms. 

116
Figure 6-9. Optical DLTS scan for UF 5% cell at a reverse bias of –0.5 V, 

λ=532 nm, and pulse width of 10 ms. 

117
Figure 6-10. The Arrhenius plot of emission rate en/T2

 vs. 1/T for DLTS scan 

at reverse bias 0.5 V, pulse height 0.7 V and pulse width 10 ms. 

118
Figure 6-11. DLTS scan for EPV cell at a reverse bias of –0.1 V, a pulse 

height of 0.3 V, and pulse width of 10 ms. 

119
ln(T2t) 

0 

-2 

-4 

-6 

-8 

-10 

-12 

0 0.0005 0.001 0.0015 0.002 0.0025 0.003 0.0035 0.004 0.0045 0.005 

Ea= 0.94 eV 

EPV CIGS 

VR = -0.1 V 

Pulse Height = 0.3 V 

Pulse Width = 10 ms 

1/T (K-1) 

Figure 6-12. The Arrhenius plot of emission rate en/T2

 vs. 1/T. 

120
 EPV_77K_C

 EPV_150K_C

 EPV_200K_C

 EPV_250K_C

 EPV_300K_C

-3 

-12 1.4x10 450x10 EPV_77K_G

 EPV_150K_G

 EPV_200K_G

 EPV_250K_G

 EPV_300K_G 

1.2 400 

1.0 

0.8 

0.6 

350 

300 

250 

Conductance (S) 

N (cm^-3) 

0.4 200 

0.2 150 

0.0 0.2 0.4 0.6 0.8 1.0 0.0 0.2 0.4 0.6 0.8 1.0 

Reverse Bias (V) Reverse Bias (V)

(a) (b) 18 

EPV_77K_C2

 EPV_150K_C2

 EPV_200K_C2

 EPV_250K_C2

 EPV_300K_C2 

50x10 4

0.2 0.3 0.4 0.5 0.6

Depletion Width (um)

 EPV_77K_N

 EPV_150K_N

 EPV_200K_N

 EPV_250K_N

 EPV_300K_N 

3 

40 2 

30 16 

10 9 

8 

7 

6 20 5 

4 

3 10 

2 

0.0 0.2 0.4 0.6 0.8 1.0 

Reverse Bias (V)

(c) (d) 

Figure 6-13. C-V data of EPV sample at different temperatures. (a) C-V (b) 

conductance-voltage (c) 1/C2

-V (d) carrier concentration-depletion width. 

1/C^2 Capacitance (F) 

121
Figure 6-14. The C-T scan for EPV cell at a reversed bias -0.1V. 

122
Figure 6-15. DLTS data for NREL sample measured at a reverse bias -0.5 V, 

pulse amplitude 0.4 V, and pulse width 10 ms. 

123
Figure 6-16. Arrhenius plot of the emission rate en/T2

 vs. 1/T. 

124
Figure 6-17. DLTS scan for the NREL sample under forward bias (0.1 V) 

injection, and heat up with reverse bias of -0.5V, pulse height of 0.7V, and 

pulse width of 10 ms. 

125
Figure 6-18. Arrhenius plot of emission rate en/T2

 vs. 1/T for the electron trap 

E1. 

126
Table 6-1. Summary of DLTS data on UF CIS sample 

VR = -0.5 V 

VH =0.4 V 

W= 10 ms 

Cool to 77 K with reverse 

bias - 0.5 V, then apply 

VR = -0.5 V VH = 0.7 V 

W= 10 ms 

Optical DLTS 

VR = -0.5 V 

λ= 532 nm 

Approximate Peak 

Temperature 

(K) 

290 290 200 150 300 

DLTS peak Sign - - + + + 

Trap Carrier Type Majority Majority Minority Minority Minority 

Trap Activation 

Energy Ea 

(eV) 

Ev +0.54 Ev +0.55 Ec -0.52 Ec - 0.16 Ec - 0.5 (?) 

Trap Density 

(cm-3) 

4.6 × 10 12 6.5 × 10 12 1.3 × 10 12 4.9 × 10 12 3.5 × 10 13 

Possible Defect 

I.D. 

CuIn CuIn VCu InCu CuIn 

Capture Cross 

Section 

σ (cm2) 

1.39 × 10 -14 

5.7 × 10 -15 1.2× 10 -18 

127
Table 6-2. Summary of DLTS and C-V data for the EPV and NREL samples. 

EPV 

CIGS 

NREL 

CIGS 

Approximate 

Peak Temperature 270 K 100K 250 K 335 K 

Activation 

Energy (eV) Ev+0.94 Ec-0.067 

DLTS peak - + + + 

Trap Carrier Type Majority Minority Minority Minority 

Na (cm-3) 3×1015 2.25×1015 

NT (cm-3) 6.5×1013 4.2×1013 

Possible defect 

I.D. 

VSe VSe 

Capture Cross 

Section σ (cm2

) 

6×10-18 

128
7 Thermodynamic Modeling of the Isothermal 500°C 

Section of the Cu-In-Se System Using a Defect Model 

7.1 Introduction 

This study aims at establishing a defect model to understand the defect chemistry of the 

Cu-In-Se system and build a database to calculate the defect and carrier concentrations. 

7.2 Experimental Phase Equilibria Data in the Cu-In-Se System 

The most complete and recent study on the phase equilibria of Cu-In-Se are conducted by 

Gödecke et al. [7-1] using DTA, XRD, TEM and SEM measurements. The isothermal section at 

500 oC is of major technical importance for the fabrication of thin films for photovoltaic devices 

because these films are often deposited at temperatures around 500 oC. A summary of all the 

phase observed equilibria is listed in Table 7-1. There are sixteen phases participating in the 

equilibria. Of which, nine are solutions: α-CIS (CuInSe2 ), βR-CIS (Cu1In3Se5), γCIS(Cu1In5Se8), βH-CIS(Cu2In4Se7), αCu, η (Cu2In1), δ (Cu7Se3), β-Cu2Se, and Ionic_liquid. The 

other phase are the line compounds In5Se7, In9Se11, In6Se7, γ- In2Se3, β-In2Se3, InSe and In4Se3. 

According to their results, in the Se-rich corner and in the In-rich corner, the isothermal section 

at 500 oC features small regions of liquid phase (L4 and L2, respectively). The α-CIS phase 

establishes equilibria with eight solid phases and two liquid phases (L4 and L2). Among the solid 

phases, αCu, η (Cu2In1), and δ (Cu7Se3) originate from the In-Cu boundary system, β-Cu2Se from 

the Cu-Se boundary system and InSe and In4Se3 from In-Se boundary system. The remaining 

two phases in equilibrium with α-CIS (CuInSe2 ) are βR-CIS (Cu1In3Se5) and βH-CIS(Cu2In4Se7). 

7.3 Thermodynamic Modeling of Defect-free Cu-In-Se System at 

773 K 

To stabilize the phase equilibria in the Cu-In-Se system with many phase interactions, it 

is important to determine the parameters when the CIS compounds are defect free, and 

assessment of this complicated Cu-In-Se ternary system begins from this point. 

7.3.1 Gibbs Energy Data of Defect-free α-CIS (CuInSe2) βR-CIS (Cu1In3Se5), γCIS(Cu1In5Se8), and βH-CIS(Cu2In4Se7) 

The sublattice model is chosen to describe the α-CIS phases. According to the sublattice 

model, the defect free Chalcopyrite CuInSe2 (stoichiometry) consists of three sublattices and is 

presented as 

(Cu)

I

1(In)

II

1(Se)

III

2 

129
The superscripts, I, II and III identify the sublattices. Analogous to α-CIS, the βR, γ, βH-CIS can 

be expressed by the following sublattices: 

(Cu)

I

1(In)

II

3(Se)

III

5 for βR-CIS 

(Cu)

I

1(In)

II

5(Se)

III

8 for γ-CIS 

(Cu)

I

2(In)

II

4(Se)

III

7 for βH-CIS 

The Gibbs energy evaluation method based on the idea that ordered defect compounds 

(ODC) are formed by repeating periodically m units of (2VCu + InCu) for every n unit of α-CIS 

(CuInSe2) [7-2, 3] are employed to evaluate the Gibbs energy of pure compound α-CIS 

(CuInSe2) βR-CIS (Cu1In3Se5), γ-CIS(Cu1In5Se8), and βH-CIS(Cu2In4Se7). The Gibbs energy 

data estimated by this method will be used as the compound Gibbs energy for the defect free α, 

βR, γ, βH-CIS so that we can express the Gibbs energy of βR, γ, βH-CIS in terms of α-CIS to 

reduce the number of parameters in the assessment when the defects are incorporated into the 

sublattices. This is because there is no experimental and calculated data available on the defect 

energy of βR, γ, βH-CIS. This method will provide an approach to express the pure compound 

and defects of Gibbs energy of βR, γ, βH-CIS in terms of the pure compound and defects Gibbs 

energy of α-CIS. In this method, the Gibbs energy data for βR-CIS (Cu1In3Se5), γCIS(Cu1In5Se8), and βH-CIS(Cu2In4Se7) are obtained by the following equations: 

°G(βR -CIS (Cu1In3Se5)) = 2.5×°G(α-CIS (CuInSe2 )+V1 (7-1) 

°G(γ-CIS(Cu1In5Se8)) = 4°G(α-CIS (CuInSe2 ))+ V2 (7-2) 

°G(βH -CIS(Cu2In4Se7)) = 3.5×°G(α-CIS (CuInSe2 ))+V3 (7-3) 

The parameters V1, V2, and V3 needed to be estimated. The estimation of V1, V2, and V3 follows 

the following procedure. First, the magnitude of V1, V2, and V3 are given by 0.5× ΔH 0 

f ,2V +In or Cu Cu 

ΔH 0 

f ,2V +In as an initial parameters since m= 0.5 or 1 for every ODC compound. Then, V1, V2, Cu Cu 

and V3 are optimized by using the Thermo-Calc Parrot module. 

7.3.2 Gibbs Energy of Other Phases 

As mentioned previously, seven other line compounds and two small liquid phase region 

also exist in this isothermal 500 °C diagram, in addition to the four CIS compounds. Among the 

solid phases, αCu, η (Cu2In1), and δ (Cu7Se3) originate from the In-Cu boundary system, β-Cu2Se 

from the Cu-Se boundary system and InSe and In4Se3 from In-Se boundary system. 

7.3.2.1 Fcc(Cu) Phase 

A regular solution model is used to describe the Gibbs energy of the fcc(Cu) phase. 

(Cu, Se, In) 

130
Combining the previous Cu-Se and Cu-In binary assessment work [7-4, 5], the ternary 

interaction parameter CuFcc 0 L ,, was added. This parameter needs to be optimized. Cu In Se 

7.3.2.2 β-Cu2Se Phase

 The β-Cu2Se phase is described by the sublattice model with two Cu sublattices and one 

Se sublattice represented by the formula: 

(Cu,In,Va)1(Se,Va)1(Cu)1 

Again a ternary interaction parameter was added to the binary Cu-Se system [7-4] to model the 

2 In solubility, β −Cu Se 0

, , : : L . Cu In Va Se Cu 

7.3.2.3 The η(Cu2In1), and δ(Cu7In3) Phases 

In the binary assessment work [7-5], both phases are described by a two sublattice model. 

To model the solubility, the two sublattice model was modified so that all the three elements will 

be involved. 

 (Cu, In, Se)2(Cu, In, Se)1 for η(Cu2In1) phase 

 (Cu, In, Se)7(Cu, In, Se)3 for δ(Cu7In3) phase 

Again a ternary interaction parameter φ L0 

Cu In Se , , :* , 

φ L0

*: , , was included, where φ stands η(Cu2In1) Cu In Se 

or δ(Cu7In3) phases. 

7.3.2.4 Liquid Phase 

The liquid phase is described by an ionic sublattice model with two sublattices, since the 

ionic two-sublattice model can describe short range ordering and is more straightforward to 

extend to higher order systems as compared to the chemical theories [7-6]. Based on the 

previous work on the modeling of the liquid phase in the Cu-Se, In-Se [7-4] and Cu-In [7-5] 

binary systems, the liquid phase is schematically described in this ternary system as 

(Cu+1,Cu+2, In+3)p(Se-2,Va-1,Se)q 

Therefore, three ternary interaction parameters were added to the ionic liquid to describe the 

Gibbs energy, +1 +2 +3 −2 +1 +2 +3 −2

Ionic _ Liq L0 

Cu ,Cu ,In :Se and Ionic _ Liq L

1 

Cu ,Cu ,In :Se . 

7.3.3 Results and Discussion 

The optimized parameters of the stable phases in this Cu-In-Se system with a defect free 

compound CIS are listed in Table 7-2. The calculated phase diagram over the entire composition 

range using the optimized parameters is shown in Figure 7-1. 

131
7.4 Defect Model of Off-Stoichiometry α-CuInSe2 

7.4.1 Modeling of Gibbs Energy 

7.4.1.1 α-CIS (CuInSe2) 

To model the homogeneity range, the defects have to be incorporated into three 

sublattices of the stoichiometric chalcopyrite compound α-CIS (CuInSe2). As reviewed in 

Chapter 2, four main defects are incorporated into this sublattice model: VCu, VSe, InCu, CuIn. 

Since the model will also incorporate the calculation of free electrons and holes, which is 

discussed in detail in the next section, it is necessary to consider the donors and acceptors in this 

phase. In the extrinsic semiconductors, the shallow-level impurities are used mainly for 

controlling the carrier concentration and the conductivity of the semiconductor, while the deeplevel impurities are used to control the recombination and hence the minority carrier lifetime in a 

semiconductor [7-7]. The Cu vacancy is a shallow level acceptor and the Se vacancy is a 

shallow-level donor. To simplify the model in this calculation, one acceptor VCu 

-1 and one donor 

VSe 

+2 are considered. Thus, the sublattices for α-CIS phase can be written as 

(Cu, In, Va, Va-1)

I

1(In, Cu)II1(Se, Va, Va+2)

III 

2 (Va)

IV (Va)V 

The superscripts, I, II and III identify the sublattices. The major constituents on the 

normal sublattices are highlighted. The IV and V sublattices are added for incorporation of 

electrons and holes. For now, just consider the first three sublattices, with the vacancy 

IV V concentration on sublattices IV and V, yVa =1 and yVa =1. 

The Gibbs energy of such a phase α can also be expressed by equation: 

Gα = ref Gα + I D . .Gα + ex Gα (7-4) 

where α stands for non-stoichiometric chalcopyrite α-CuInSe2 phase. Now, the reference state 

can be expressed in the form 

132
ref Gα I II III IV V o α I II III IV V o α = y y y yVa yVa G + y y y yVa yVa G 

Cu In Se Cu : In : Se : Va : Va In In Se In : In : Se : Va : Va 

I II III IV V o α I II III IV V o α + y y y yVa yVa G + y −1 y y yVa yVa G −1

Va In Se Va : In : Se : Va : Va Va In Se Va : In : Se : Va : Va 

I II III IV V o α I II III IV V o α + y y y y y G Va Va + y y y y y G Va Va Cu Cu Se Cu : Cu : Se : Va : Va In Cu Se In : Cu : Se : Va : Va 

I II III IV V o α I II III IV V o α + y y y yVa yVa G + y −1 y y yVa yVa G −1

Va Cu Se Va : Cu : Se : Va : Va Va Cu Se Va : Cu : Se : Va : Va 

I II III IV V o α I II III IV V o α + y y y y y G Va Va + y y y y y G Va Va Cu In Va Cu : In : Va : Va : Va In In Va In : In : Va : Va : Va 

I II III IV V o α I II III IV V o α + y y y yVa yVa G + y −1 y y yVa yVa G −1

Va In Va Va : In : Va : Va : Va Va In Va Va : In : Va : Va : Va 

I II III IV V o α I II III IV V o α + y y y y y G Va Va + y y y y y G Va Va Cu Cu Va Cu : Cu : Va : Va : Va In Cu Va In : Cu : Va : Va : Va 

I II III IV V o α I II III IV V o α + y y y y y G Va Va + y y y y y G −1 Va Va −1

Va Cu Va Va : Cu : Va : Va : Va Va Cu Va Va : Cu : Va : Va : Va 

I II III IV V o α I II III IV V o α + y yy +2 y y Va Va G +2 + y y y +2 yVa yVa G +2

Cu In Va Cu : In : Va : Va : Va In In Va In : In : Va : Va : Va 

I II III IV V o α I II III IV V o α + y y y +2 yVa yVa G +2 + y −1 y y +2 yVa yVa G −1 +2

Va In Va Va : In : Va : Va : Va Va In Va Va : In : Va : Va : Va 

I II III IV V o α I II III IV V o α + yyy +2 y y G +2 + y y y +2 y y G Va Va Va Va +2

Cu Cu Va Cu : Cu : Va : Va : Va In Cu Va In : Cu : Va : Va : Va

I II III IV V o α I II III IV V o

Gα (7-5) 

+ yyy +2 yVa yVa G +2 + y −1 y y yVa yVa −1 +2

Va Cu Va Va : Cu : Va : Va : Va Va Cu Va Va : Cu : Va : Va : Va 

Including the ideal entropy of mixing 

IDGα = + RT( yI ln yI + yI ln yI + yI ln yI + yI 

−1 ln yI −1 ) Cu Cu In In Va Va Va Va 

+ RT( y

II ln yII + yII ln yII ) Cu Cu In In (7-6) 

+ RT( y

III ln yIII + yIII ln yIII + yIII 

+2 ln y

III 

+2 ) Se Se Va Va Va Va 

and finding the excess Gibbs energy of mixing 

133
⎡y

II 

y

III 0L

α + y

II 

y

III 0L

α ⎤ ⎢ ( ) ( ) ⎥ EX α I I Cu Se Cu In Cu Se Va Va , : : : : In Se Cu In In Se Va Va , : : : :

G = y y ⎢ ⎥ Cu In ⎢+y

II 

y

III 0L

α + y

II 

y

III 0L

α ⎥ ⎢ Cu Va ( , : : : : ) In Va ( , : : : : ) ⎣ Cu In Cu Va Va Va Cu In In Va Va Va ⎥⎦

⎡ II III 0 α II III 0 α ⎤ y y ( L ) + yy ( L ) I I ⎢ Cu Se Cu Va Cu Se Va , : : : :Va In Se Cu Va In Se Va Va , : : : : ⎥

+ y y Cu Va ⎢

⎢+y

II 

y

III 0L

α + y

II 

y

III 0L

α

⎥

⎥ ⎢ Cu Va ( , : : : : ) In Va ( , : : : : ) ⎣ Cu Va Cu Va Va Va Cu Va In Va Va Va ⎦⎥

⎡ II III ⎛ 0 α ⎞ II III ⎛ 0 α ⎞ ⎤

I I ⎢ y y ⎝

⎜ L , −1 : : : : ⎠

⎟+ y

In ySe ⎝

⎜ L

Cu Va , −1 : In : Se :Va Va ⎠ Cu Se ⎟ ⎥ Cu Va Cu Se Va Va : 

+ y ⎢ ⎥ Cu −1 y

Va ⎢

+y

II 

y

III ⎛ 0L

α ⎞ + y

II 

y

III ⎛ 0L

α ⎞ ⎥

⎣ ⎝ , : : : : ⎝ Cu Va , : : : : ⎦ ⎢ Cu Va ⎜ Cu Va −1 Cu Va Va Va ⎟

⎠ In Va ⎜ −1 In Va Va Va ⎟⎠ ⎥

⎡ y

I 

y

III ( 0L

α ) + y

I 

y

III ( 0L

α ) ⎤ ⎢ Cu Se Cu Cu In : , : Se Va Va : : In Se In Cu In Se Va Va : , : : : ⎥

II II ⎢ + I III 0 α II III 0 α ⎥ y y 

⎢

+y y ( L ) + y y ( L ) ⎥

Cu In ⎢ Cu Va Cu Cu In Va Va Va : , : : : In Va In Cu In Va Va Va : , : : : ⎥ ⎢ I III 0 α I III 0 α ⎥

⎢⎣ Va Se ( Va Cu In : , : Se Va Va : : ) Va Va ( : , : : : ) ⎢+y y L + y y LVa Cu In Va Va Va ⎥

⎥⎦

⎡ I II 0 α I II 0 α ⎤ ⎢ Cu Cu ( : : , : : ) In Cu ( : : Se Va Va Va : : )⎥ y y LCu Cu Se Va Va Va + y y LIn Cu , 

+ y

III III ⎢+y

I 

y

II 0L

α + y

I 

y

II 0L

α ⎥

Va ySe ⎢

⎢

Cu In ( : : , : : ) In In ( : : , : : ) ⎥

⎥

Cu In Se Va Va Va In In Se Va Va Va 

⎢ I II 0 α I II 0 α ⎥ ⎢+y y L + y y L ⎥ ⎢⎣ Va Cu ( , ) ( Va In : Se Va Va Va ) Va Cu ⎥⎦ : : Se Va Va In : , : : 

⎡ I II ⎛ 0 α ⎞ I II ⎛ 0 α ⎞⎤

y y ⎜ L +2 ⎟+ yy ⎜ L ⎢ +2 ⎟ Cu Cu Cu Cu Se Va : : , : : In : Cu : Se Va , : : ⎥ ⎢ ⎝ Va Va ⎠ In Cu ⎝ Va Va ⎠⎥

+ y

III +2 III ⎢

⎢

+y

I 

y

II ( 0L

α ) + y

I 

y

II ⎜

⎛ 0L

α +2 ⎟

⎞ ⎥

⎥

Va y

Se ⎢ Cu In : : , : : In In ⎝ , ⎠ ⎥ Cu In Se Va Va Va In In Se Va : :

⎢ I II ⎛ 0 α ⎞ I II ⎛ 0 α ⎞⎥

+yy L + y y L ⎣ ⎝ Va Cu : : Se Va , :Va Va ⎝ : : , : : ⎦ ⎢ Va Cu ⎜ +2 : ⎟

⎠ Va In ⎜ Va In Se Va +2 Va Va ⎟⎠

⎥

(7-7) 

where yi

I 

, yi

II and yiIII are site fractions on the first three sublattices. oG is the Gibbs energy of a 

real or fictitious compound. They are called end-members and there are 24 of them as listed in 

the above equations. The 0

L parameters are the interaction parameter between the end members. 

The Gibbs energy of end-members, which have one defect, 

o o o o o ( Gα , Gα , Gα , Gα , Gα −1

Cu : Cu : Se : Va : Va Va : In : Se : Va : Va In : In : Se : Va : Va Cu : In : Va : Va : Va Va : In : Se : Va : Va 

, 

o

Gα , etc) is obtained from defect formation energy and entropy calculation in the +2

Cu : In : Va : Va : Va 

literature, which are based on the deviation from that of the same quantity of non-defective CIS. 

For the end-members which have two defects, a simple reciprocal relation [7-8] is used to 

obtain its Gibbs energy. 

o

Gα =oGα + oGα -

o

Gα

Va : Cu : Se : Va : Va Cu : Cu : Se : Va : Va Va : In : Se : Va : Va Cu : In : Se : Va : Va 

134
o

Gα −1 =oGα + oGα −1 - oGα

Va : Cu : Se : Va : Va Cu : Cu : Se : Va : Va Va : In : Se : Va : Va Cu : In : Se : Va : Va

o

Gα =oGα + oGα - oGα

In : Cu : Se : Va : Va Cu : Cu : Se : Va : Va In : In : Se : Va : Va Cu : In : Se : Va : Va 

o

Gα =oGα + oGα - oGα

Va : In : Va : Va : Va Cu : In : Va : Va : Va Va : In : Se : Va : Va Cu : In : Se : Va : Va 

o o o o Gα −1 = Gα + Gα −1 - Gα

Va : In : Va : Va : Va Cu : In : Va : Va : Va Va : In : Se : Va : Va Cu : In : Se : Va : Va 

o

Gα =oGα + oGα - oGα

In : In : Va : Va : Va Cu : In : Va : Va : Va In : In : Se : Va : Va Cu : In : Se : Va : Va 

o

Gα =oGα + oGα - oGα +2 +2

Va : In : Va : Va : Va Cu : In : Va : Va : Va Va : In : Se : Va : Va Cu : In : Se : Va : Va 

o

Gα =oGα + oGα - oGα −1 +2 +2 −1

Va : In : Va : Va : Va Cu : In : Va : Va : Va Va : In : Se : Va : Va Cu : In : Se : Va : Va 

o

Gα =oGα + oGα - oGα +2 +2

In : In : Va : Va : Va Cu : In : Va : Va : Va In : In : Se : Va : Va Cu : In : Se : Va : Va 

o

Gα =oGα + oGα - oGα

Cu : Cu : Va : Va : Va Cu : Cu : Se : Va : Va Cu : In : Va : Va : Va Cu : In : Se : Va : Va 

o

Gα =oGα + oGα - oGα +2 +2 (7-8) Cu : Cu : Va : Va : Va Cu : Cu : Se : Va : Va Cu : In : Va : Va : Va Cu : In : Se : Va : Va 

The Gibbs energy of end-members, which have three or more than three defects are 

simply put to an arbitrarily large value [7-8]. 

The interaction parameter α L0

 are assumed to follow relationships that reduce the 

parameters; 

135
 

 

 

α 0 α 0 α 0 α 0 LCu In , :*:* V a V a : = L , : : : : = LCu In Cu Se Va Va , : : : : = L Cu In In Se Va Va Cu In In Va Va Va , : : : :

Cu In Cu Va Va Va Va Va ,: V a V a = α L0

, : : : : = α L0 

Cu In In Va ,:: +2 : : = α L0 Cu In Cu Va : + 2 : : 

α 0 α 0 α 0 α 0 LCu Va , *:*: V a V a = L , : : : : = LCu Va Cu Se Va Va : : : : = L : Cu Va I n S e V a V a , C u V a In V a V a V a , : : : :

= α L0 = α L0 = α L0 

Cu Va Cu Va Va Va , : : : Cu Va In Va , : : + 2 :V a V a : Cu Va Cu Va : : + 2 : : , V a V a : 

α 0 α 0 α 0 α 0 L*: , V a V a : = L : , : : : = LIn C u : , In Se Va V : : : a = L Cu In :*: Cu Cu In Se Va Va In Cu In Va Va Va : , : : :

= α L0

: , : : : = α L0 

+ 2 = α L0 + 2 (7-9) C u Cu In Va Va Va C u C u In Va : , : : : Va Va C u C u In Va : , : : : Va Va 

= α L0 = α L0 = α L0 

: , : : : , : : Va C u In Va , : + 2 Va Cu In Se Va Va : : : Va C u In Va Va Va : : V a V a 

α 0 α 0 α 0 α 0 L*:* : , : : = LC u C u Se Va Va Va : : , : : = L Se Va Va Va In Cu Se Va Va Va : : , : : = LVa Cu Se Va Va Va : : , : :

α 0 α 0 α 0 = L : : , : : = L Cu In Se Va Va Va In In Se Va Va Va : : , : : = LVa In Se Va Va Va :: , : :

α 0 α 0 α 0 α 0 L = L = L = L *:*: , +2 V a V a : : : , + 2 :Va Va : : , +2 V a V a : : : , + 2 Se Va : Cu Cu Se Va : In Cu Se Va : Va Cu Se Va :V a V a : 

= α L0 = α L0 = α L0 

Cu In Se Va : : , + 2 : : In In Se Va , +2 : : Va In Se Va , +2 V a V a :: V a V a :: : : V a V a 

7.4.1.2 II. βR-CIS (Cu1In3Se5), γ-CIS(Cu1In5Se8), and βH -CIS(Cu2In4Se7) 

Similar to the modeling the non-stoichiometric α-CIS (CuInSe2 ) phase, a three sublattice 

model was also used to describe these compounds. Since the electrons and holes are not 

modeled in these CIS phases, sublattices IV and V are not added. 

For βR -CIS (Cu1In3Se5) phase: (Cu, Va, In)I

1(In, Cu)II3(Se, Va)III5 

For γ - CIS(Cu1In5Se8) phase: (Cu, Va, In)I

1(In, Cu)II5(Se, Va)III8 

For βH-CIS(Cu2In4Se7) phase : (Cu, Va, In)I

2(In, Cu)II4(Se, Va)III7 

As mentioned previously, these ODC compounds are formed by repeating periodically m −1 +2 units of 2VCu + In Cu for every n unit of α-CIS (CuInSe2) [7-9, 10]. The Gibbs energy for these 

compounds can also be expressed in the similar way as α-CIS (CuInSe2). The value of Gibbs 

energy of the end-members of one or two defects in these compounds in relation to the Gibbs 

energy of the end-members of α-CIS (CuInSe2) can be set by the following expression: 

o

Gβ R = 2.5×o

Gα , i : j : k : Va : Va i : j : k : Va : Va 

o

Gβ H = 3.5×o

Gα , i : j : k : Va : Va i : j : k : Va : Va 

o

Gγ = 4×o

Gα (7-10) i : j : k : Va : Va i : j : k : Va : Va 

136
where i, j, and k represent the various sublattice constituents. For example, o

Gβ R = 

Cu : Cu : Se : Va : Va 

o 2.5× Gα . Cu : Cu : Se : Va : Va 

7.4.2 Calculation of the Carrier Concentration in α-CIS (CuInSe2 ) 

The free electrons and holes concentration can be calculated if they are also included in the 

thermodynamic description. They are supposed to be distributed on different numbers of sites 

than there are sites in the sublattices [7-11]. As known from the Fermi-Dirac distribution 

function for electrons in the conduction band and holes in the valence band Nc and Nv stand for 

the effective density of the conduction and valence band states, respectively. So, let Nc and Nv 

be the number of sites that electrons and holes can occupy per formula unit. 

According to the above description, the α-CIS phase can be described by the following 

sublattices. 

(Cu, Va, Va-1, In)I

1(In, Cu)II1(Se, Va, Va+2)

III 

2(Va, e-1)

IV

Nc(Va, h+1)

V

Nv 

This model can be directly described with the Compound Energy Model (CEM). 

However, CEM requires Nc and Nv to be constant. Chen et al. [7-8] provided an alternative way 

to solve this problem and transferred the sublattice description to the following: 

(Cu, Va, In)I

1(In, Va, Cu)II1(Se, Va, Va+2)

III 

2(Va, e-1)

IV

1(Va, h+1)

V

1 

Here is briefly described their transformation methodology by applying it to the α-CIS 

phase. The effect of adding the IV sublattice for electrons is, instead of having the 

I II III 0 term, y y y G , the equation * * * *:*:* 

I II III IV 0 I II III IV 0 I II III IV 0 I II III IV 0 y y y y G + y y y y 1 G −1 = y y y (1 − y −1 )G + y y y y −1 G * * * Va *:*:*: Va * * * − e− *:*:*: e * * * e *:*:*: Va * * * e *:*:*: e 1 

I II III 0 0 0 I II III IV = yy y G + (G − −G ) yy y y * * * *:*:*: Va 1 *:*:*: e 1 *:*:*: Va * * * e−

( yIV + IV =1) (7-11) Va ye−1 

0 0 For simplicity, it is assumed all values (G −1 −G*:*:*: Va ) are equal. However, the quantities refer *:*:*: e 

to Nc moles of electron. Denote the electron chemical potential as 

0 0,ref 0 0 ΔG -1 -G = (G −1 −G*:*:*: Va ) / Nc (7-12) e *:*:*: e 

The G0,ref can be immediately omitted due to the condition of electroneutrality. It will drop out 

of the final expression for G if the same reference is used for all charged constituents. 

Substituting equation (7-12) into equation (7-11) yields 

137
I II III IV 0 I II III IV 0 y y y y G * * * Va *:*:*:Va + y y y y G * * * e−1 *:*:*:e−1 

(7-13) I II III 0 0 I II III IV = y y y G + ΔG −1 y y y (N • y −1 ) * * * *:*:*:Va e * * * c e 

I II III 0 Since y y y =1, all the values of ΔG ∑∑∑ −1 are equal and similar expressions for the * * * e I II III 

holes exist 

GNew =GOld 

0 IV I II III 0 IV I II III 0 IV IV + ΔGN −1 ( • y ) y y y G +Δ (N • y ) y y y +ΔG (N • y )(N • y ) c −1 ∑∑∑ * * * +1 V +1 ∑∑∑ * * * −1 +1 c −1 V +1 e e h h e h e h I II III I II III 

c[ e

IV 

−1 ln y

e

IV 

−1 +(1− y

e

IV 

−1 )ln(1− y

e

IV 

−1 )]+RTN V[y

h

IV 

+1 l y

h

IV 

+1 + − y

h

IV 

+1 )ln(1− y

h

IV +RTN y n (1 +1 )]

=GOld 

0 IV I II III (7-14) [ G −RT ln(Nc + Δ )](N • y )∑∑∑ * * −1 c −1 y y y * + e e I II III

0 IV I II III 0 IV IV [ΔG RT +1 − ln(N )](N • y ) y y y +Δ (N • y )(Ny ) v V +1 ∑∑∑ * * * G −1 +1 c −1 V • +1 h suuuuuuuuuuur h e h e h I II III 

+RT{(N y • IV )ln(N yIV + N •(1− y

IV )]ln[N •(1− IV 

c −1 c • −1 ) [ c −1 c y −1 )}+ e e e e 

RT{(N • y

IV )ln(N • yIV + N •(1− yIV )ln[N •(1 IV ) [ − y )]} V h+1 surv h+1 v h+1 surv h+1 

Based on the above formula, the full α-phase CIS is represented as: 

(Cu, Va, In)I

1(In, Va, Cu)II1(Se, Va)III2(Va, e-1)

IV

1(Va, h+1)

V

1 

Defining the midgap energy as the energy reference, the following relations exist 

0 0 G*:*:*:Va Va : = G*:*:*

E 0 0 g G*:*:*: −1

: e Va = G*:*:* + − RT ln NC 2 (7-15) E 0 0 g G*:*:*: : +1 = G*:*:* +

2 Va h − RT ln NV

0 0 G*:*:*: −1

: eh +1 = G*:*:* + Eg − RT ln NC − RT ln NV

The last relation is not meant to be realistic, but it has a negligible effect because the 

contributions of both free electrons and holes are small relative to site density. The effective 

densities Nc and Nv are calculated by the following equations [7-8]: 

* 2 3/ 2 2 N = 2(2πm kT / h ) • (a c / 8) c c (7-16) * 2 3/ 2 2 N = 2(2πm kT / h ) • (a c / 8) v v 

(ac / 8) is the volume per mole of the formula units and is added to get the densities counted 

per mole of formula units. mc

* and mv* are the effective masses of electrons at the conduction 

138

2 
band edge and holes at the valence band edge, respectively. k is Boltzman’s constant, and h is 

Planck’s constant. The values of all the parameters are summarized in Table 7-3 [7-12, 13]. 

7.4.3 Results and Discussion 

A reasonable agreement for the α-CIS phase was obtained between the model-calculated 

values and the phase equilibrium available in the literature as shown in Figure 7-2. The 

optimized parameters are listed in the Table 7-6. The calculated defect and carrier concentration 

data is shown in Figure 7-3. 

Further work needs to be done to get better agreement of phase equilibrium for all the 

phases. 

7.5 References cited 

[7-1] T. Gödecke, T. Haalboom, and F. Ernst , Zeitschrift für Metallkunde, Vol. 91, No. 8 

pp.622-662, 2000. 

[7-2] R.D. Wieting, the 29th IEEE Photovoltaic Specialist Conference, 2002. 

[7-3] D.X. Liao, A. Rockett, Applied Physics Letters, Vol. 82, No. 17, pp. 2829-2831, 2003. 

[7-4] C-H Chang, Processing and Characterization of Copper Indium Selenide for 

Photovoltaic Applications, Ph.D. Dissertation, University of Florida, 2000. 

[7-5] P.R. Subramanian and D.E. Laughlin. Bulletin of Alloy Phase Diagrams 10, p. 554, 1989. 

[7-6] W. Zhuang, C-H Chang, J. Sheng, and T.J. Anderson, CALPHAD, 2002. 

[7-7] S.S. Li, Semiconductor Physical Electronics, Plenum Press, New York and London, 

1993. 

[7-8] Q. Chen, M. Hillert, B. Sundman, W.A. Oates, S. G. Fries, and R. Schmid-Fetzer, 

Journal of Electronic Materials, Vol. 27, No. 3 pp.961-971, 1998. 

[7-9] S. B. Zhang, Su-Huai Wei, and Alex Zunger, Physical Review B, Vol. 57, No. 16 pp9642 

–9656, 1998. 

[7-10] S. B. Zhang, Su-Huai Wei, and Alex Zunger, Physical Review Letters, Vol. 78, No. 21 

pp4059 – 4062, 1997. 

[7-11] T. J. Anderson, CALPHAD, Vol. 21, No. 2, pp. 266-282, 1997. 

[7-12] John B. Mooney and Robert H. Lamoreaux, Solar Cells, Vol. 16. pp. 211-220, 1986. 

[7-13] O. Madelung, Numerical Data and Functional Relationships in Science and Technology 

Semiconductors: Physics of III-V Compounds, Berlin, New York : Springer-Verlag, 1991. 

139
[7-14] P.R. Subramanian and D.E. Laughlin. Bulletin of Alloy Phase Diagrams 10, p. 554, 1989. 

[7-15] M. A. Contreras, H. Wiesner, D. Niles, K. Ramanathan, R. Maston, J. Tuttle, J. Keane, 

and R. Noufi, the 25th IEEE Photovoltaic Specialist Conference, Washington, D.C. 1996. 

[7-16] A. L. Fahrenbruch, and R. H. Bube, Fundamentals of Solar Cells, Academic Press, New 

York-London, 1983. 

[7-17] J.C. Mikkelsen, Journal of Electronic Materials, Vol. 10, pp. 541, 1981. 

140
Figure 7-1. Simulation results of defect free Cu-In-Se system. 

141
Figure 7-2. Predicted phase diagram of the Cu-In-Se system at 773 K that 

indicates non stoichiometric CuInSe2. 

142
1.00E+22 -1

Defect or Carrier Concentration (1/cm3) 1.00E+20 

1.00E+18 

1.00E+16 

1.00E+14 

1.00E+12 

1.00E+10 

1.00E+08 

1.00E+06 

1.00E+04 

1.00E+02 

1.00E+00 

1.00E-02 

1.00E-04 

1.00E-06 

1.00E-08 

VCu 

VSe

+2 

VSe 

h+1 

CuIn 

-1 e

VCu 

InCu 

0.0005 0.001 0.0015 0.002 0.0025 0.003 0.0035 

1/T (K) 

Figure 7-3. Calculated defect and carrier concentration at the α-CIS/β-Cu2Se 

phase boundary. 

143
Table 7-1. Experimental compositions in the Cu-In-Se system at 773K [7-1]. 

Phases Composition, 

at. % Se 

Composition, 

at. % Cu 

α-CIS/βRCIS/IONIC_LIQ 

0.5294/0.553/0.9852 0.1963/0.149/0.00142 

IONIC_LIQ/α-CIS 

0.5234/0.9904 

0.5195/0.9926 

0.5141/0.99285 

0.2052/0.001896 

0.2152/0.001942 

0.2250/0.003625 

α-CIS /β-Cu2Se 

/IONIC_LIQ 

0.5096/0.3794/0.994 0.2354/0.6157/0.00592 

α-CIS /β-Cu2Se 

/CU_FCC 

0.5017/0.33646/0.0054 0.23895/0.6562/0.9061 

α-CIS /CU_FCC 0.5003/0/0.00296 0.2371/0.7/0.904 

α-CIS /δ (Cu7Se3) 

/η(Cu2In1) 

0.4988/0/0.0064 0.2353/0.7/0.6388 

α-CIS / 

IN4SE3/η(Cu2In1) 

0.4983/0.4286/0.001 0.2326/0/0.6225 

α-CIS 

/IN1SE1_A/IN4SE3 

0.5012/0.5/0.4286 0.2278/0/0 

α-CIS / IN1SE1_A/βHCIS 

0.5052/0.5/0.5141 0.22156/0/0.1388 

α-CIS /βR-CIS /βH-CIS 0.5101/0.52795/0.523 0.2123/0.1649/0.161 

IONIC_LIQ/γ- In2Se3/γCIS 

0.961/0.6/0.5942 0/0/0.0332 

IONIC_LIQ/γ-CIS /βRCIS 

0.978/0.5932/0.589 0.02/0.0557/0.095 

In6Se7/γ-CIS//βR-CIS 0.538/0.5675/0.5537 0/0.0592/0.088 

In6Se7/γ-CIS/ In9Se11 0.538/0.5720/0.55 0/0.052/0 

In5Se7/γ-CIS/ In9Se11 0.5933/0.5744/0.55 0/0.0465/0 

In5Se7/γ-CIS/β-In2Se3 0.5933/0.5819/0.6 0/0.0368/0 

144
Table 7-2. Summary of the optimized parameters. Note: all the units are in J/mol. 

Phases and 

Functions 

 Parameters 

α-CIS 

(CuInSe2) 

Cu In Se Va Va : : : :

G

o α = CIS_A 

βR-CIS 

(Cu1In3Se5) 

: : : : 

R 

Cu In Se Va Va 

G

o β = 2.5×CIS_A+V1 

βH-CIS 

(Cu2In4Se7) : : : : 

H 

Cu In Se Va Va 

G

o β = 3.5×CIS_A+V2 

γ-CIS 

(Cu1In5Se8) Cu In Se Va Va : : : :

G

o γ = 4×CIS_A+V3 

η (Cu2In1) In addition to all the parameters from Cu-In binary system [7-14], the 

following ternary interaction parameters are added: 

0 

,, 

η LCu In Se:* =-1029637 0 

*: , , 

η L Cu In Se =-1029637 

β-Cu2Se In addition to all the parameters from Cu-Se binary system [7-4], the 

following ternary interaction parameter is added: 

2 0 

, , : : 

Cu SeLCu In Va Se Cu 

β − =-20000 

δ (Cu7In3) In addition to all the parameters from Cu-In binary system [7-14], the 

following ternary interaction parameters are added 

0 

,, 

δ LCu In Se:* =-90000 0 

*: , , 

δ L Cu In Se =-90000 

Fcc (Cu) 

In addition to all the parameters from Cu-In and Cu-Se binary 

system[44, 50], the following ternary interaction parameter is added 

0 

,, 

CuFcc LCu In Se =-90000 

Ionic_Liq In addition to all the parameters from Cu-In In-Se and Cu-Se binary[7­

15, 16] system, the following ternary interaction parameters are added: 

1 2 3 2 

_ 0 

, , : 

Ionic Liq 

Cu Cu In Se 

L + + + − =-46015; 1 2 3 2 

_ 1 

, , : 

Ionic Liq 

Cu Cu In Se 

L + + + − =9015 

145
Table 7-3. Summary of some of α-CIS property values. 

Parameter Value Ref 

A 5.782 Å [7-13] 

C 2.0097 a = 11.6201 Å [7-13] 

E 

g

0.965 -1.086×10-4 T2

/(T+97) (eV) [7-17] 

* 

m c 0.091 m0 [7-17] 

* 

m v 0.091 m0 ( = 9.11×10-31 kg) [7-17] 

k 8.617 × 10-5 eV/K [7-17] 

h 6.63 × 10-34 J ⋅ S [7-17] 

Nc ≈ Nv 6.34 × 10-13 T3/2 [7-17] 

146
Table 7-4. Summary of the optimized parameters for off-stoichiometry CIS. 

Phases 

and Functions Parameters 

α-CIS o

Gα = CIS_A (CuInSe Cu : In : Se : Va : Va 2) 

o

Gα = VIS_A Va : In : Se : Va : Va 

o

Gα = EIS_A −1

Va : In : Se : Va : Va 

o

Gα = IIS_A In : In : Se : Va : Va 

o

Gα = CS_A Cu : Cu : Se : Va : Va 

o

Gα = CIV_A Cu : In : Va : Va : Va 

o

Gα = CIH_A +2

Cu : In : Va : Va : Va 

o

Gα = CCS_A+VIS_A-CIS_A Va : Cu : Se : Va : Va 

o

Gα = CCS_A+EIS_A-CIS_A −1

Va : Cu : Se : Va : Va 

o

Gα = CCS_A+IIS_A-CIS_A In : Cu : Se : Va : Va 

o

Gα = VIS_A+CIV_A-CIS_A Va : In : Va : Va : Va 

o

Gα = EIS_A+CIV_A-CIS_A −1

Va : In : Va : Va : Va 

o

Gα = IIS_A+CIV_A-CIS_A In Va Va Va : In : : :

o

Gα = CIH_A+VIS_A-CIS_A +2

Va : In : Va : Va : Va 

o

Gα = CIH_A+EIS_A-CIS_A −1 +2

Va : In : Va : Va : Va 

o

Gα = CIH_A+IIS_A-CIS_A +2

In : In : Va : Va : Va 

o

Gα = CIV_A+CCS_A-CIS_A Cu : Cu : Va : Va : Va 

o

Gα = CCS_A+CIH_A-CIS_A +2

Cu : Cu : Va : Va : Va 

o

Gα = CIS_A+TEGN −1

Cu : In : Se : Va : Va 

147
o

Gα = VIS_A+TEGN −1

Va : In : Se : Va : Va 

o

Gα = EIS_A+TEGN −1 −1

Va : In : Se : Va : Va 

o

Gα = IIS_A+TEGN −1

In : In : Se : Va : Va 

o

Gα = CCS_A+TEGN −1

Cu : Cu : Se : Va : Va 

o

Gα = CIV_A+TEGN −1

Cu : In : Va : Va : Va 

o

Gα = CIH_A+TEGN +2 −1

Cu : In : Va : Va : Va 

o

Gα = CIS_A+TEGH +1

Cu : In : Se : Va : Va 

o

Gα = VIS_A+TEGH +1

Va : In : Se : Va : Va 

o

Gα = EIS_A+TEGH −1 +1

Va : In : Se : Va : Va 

o

Gα = IIS_A+TEGH +1

In : In : Se : Va : Va 

o

Gα = CCS_A+TEGH +1

Cu : Cu : Se : Va : Va 

o

Gα = CIV_A+TEGH +1

Cu : In : Va : Va : Va 

o

Gα = CIH_A+TEGH +2 +1

Cu : In : Va : Va : Va 

o

Gα = CIS_A+TEGH+TEGN −1 +1

Cu : In : Se : Va : Va 

α L0 

Cu In , :*:*: : = LCI Va Va 

α L0 = LCV Cu Va , :*:*: : Va Va 

α L0

, −1

:*:*: Va Va 

= LCE Cu Va : 

148
α L0

*: , Va Va : = LIC Cu In :*: 

α L0 = LSV *:*: , : : Se Va Va Va 

α L0

*:*: , +2 Va Va : = LSH Se Va : 

βR-CIS 

(Cu1In3Se5) 

o

Gβ R = 2.5×CIS_A+44500 Cu : In : Se : Va : Va 

o

Gβ R = 2.5×CCS_A Cu : Cu : Se : Va : Va 

o

Gβ R = 2.5×VIS_A Va : In : Se : Va : Va 

o

Gβ R = 2.5×IIS_A In : In : Se : Va : Va 

o

Gβ R = 2.5×CIV_A Cu : In : Va : Va : Va 

o

Gβ R = 2.5×EIS_A −1

Va : In : Se : Va : Va 

o

Gβ R = 2.5×CIH_A +2

Cu : In : Va : Va : Va 

o

Gβ R = 2.5×CCS_A+2.5×VIS_A-2.5×CIS_A Va : Cu : Se : Va : Va 

o

Gβ R = 2.5×CCS_A+2.5×EIS_A-2.5×CIS_A −1

Va : Cu : Se : Va : Va 

o

Gβ R = 2.5×CCS_A+2.5×IIS_A-2.5×CIS_A In : Cu : Se : Va : Va 

o

Gβ R = 2.5×VIS_A+2.5×CIV_A-2.5×CIS_A Va : In : Va : Va : Va 

o

Gβ R = 2.5×EIS_A+2.5×CIV_A-2.5×CIS_A −1

Va : In : Va : Va : Va 

o

Gβ R = 2.5× IIS_A+2.5×CIV_A-2.5×CIS_A In : In : Va : Va : Va 

o

Gβ R = 2.5×CIH_A+2.5×VIS_A-2.5×CIS_A +2

Va : In : Va : Va : Va 

o

Gβ R = 2.5×CIH_A+2.5×EIS_A-2.5×CIS_A −1 +2

Va : In : Va : Va : Va 

149
o

Gβ R = 2.5×CIH_A+2.5×IIS_A-2.5×CIS_A +2

In : In : Va : Va : Va 

o

Gβ R = 2.5×CIV_A+2.5×CCS_A-2.5×CIS_A Cu : Cu : Va : Va : Va 

o

Gβ R = 2.5×CCS_A+2.5×CIH_A-2.5×CIS_A +2

Cu : Cu : Va : Va : Va 

o

Gβ R = 2.5×CIS_A+TEGN −1

Cu : In : Se : Va : Va 

o

Gβ R = 2.5×VIS_A+TEGN −1

Va : In : Se : Va : Va 

o

Gβ R = 2.5×EIS_A+TEGN −1 −1

Va : In : Se : Va : Va 

o

Gβ R = 2.5×IIS_A+TEGN −1

In : In : Se : Va : Va 

o

Gβ R = 2.5×CCS_A+TEGN −1

Cu : Cu : Se : Va : Va 

o

Gβ R = 2.5×CIV_A+TEGN −1

Cu : In : Va : Va : Va 

o

Gβ R = 2.5×CIH_A+TEGN +2 −1

Cu : In : Va : Va : Va 

o

Gβ R = 2.5×CIS_A+TEGH +1

Cu : In : Se : Va : Va 

o

Gβ R = 2.5×VIS_A+TEGH +1

Va : In : Se : Va : Va 

o

Gβ R = 2.5×EIS_A+TEGH −1 +1

Va : In : Se : Va : Va 

o

Gβ R = 2.5×IIS_A+TEGH +1

In : In : Se : Va : Va 

o

Gβ R = 2.5×CCS_A+TEGH +1

Cu : Cu : Se : Va : Va 

o

Gβ R = 2.5×CIV_A+TEGH +1

Cu : In : Va : Va : Va 

o

Gβ R = 2.5×CIH_A+TEGH +2 +1

Cu : In : Va : Va : Va 

o

Gβ R = 2.5×CIS_A+TEGH+TEGN −1 +1

Cu : In : Se : Va : Va 

βH-CIS o

Gβ H = 3.5×CIS_A+44500 (Cu Cu : In : Se : Va : Va 2In4Se7) 

o

Gβ H = 3.5×CCS_A Cu : Cu : Se : Va : Va 

o

Gβ H = 3.5×VIS_A Va : In : Se : Va : Va 

150
o

Gβ H = 3.5×IIS_A In : In : Se : Va : Va 

o

Gβ H = 3.5×CIV_A Cu : In : Va : Va : Va 

o

Gβ H = 3.5×EIS_A −1

Va : In : Se : Va : Va 

o

Gβ H = 3.5×CIH_A +2

Cu : In : Va : Va : Va 

o

Gβ H = 3.5×CCS_A+3.5×VIS_A-3.5×CIS_A Va : Cu : Se : Va : Va 

o

Gβ H = 3.5×CCS_A+3.5×EIS_A-3.5×CIS_A −1

Va : Cu : Se : Va : Va 

o

Gβ H = 3.5×CCS_A+3.5×IIS_A-3.5×CIS_A In : Cu : Se : Va : Va 

o

Gβ H = 3.5×VIS_A+3.5×CIV_A-3.5×CIS_A Va : In : Va : Va : Va 

o

Gβ H = 3.5×EIS_A+3.5×CIV_A-3.5×CIS_A −1

Va : In : Va : Va : Va 

o

Gβ H = 3.5× IIS_A+3.5×CIV_A-3.5×CIS_A In : In : Va : Va : Va 

o

Gβ H = 3.5×CIH_A+3.5×VIS_A-3.5×CIS_A +2

Va : In : Va : Va : Va 

o

Gβ H = 3.5×CIH_A+3.5×EIS_A-3.5×CIS_A −1 +2

Va : In : Va : Va : Va 

o

Gβ H = 3.5×CIH_A+3.5×IIS_A-3.5×CIS_A +2

In : In : Va : Va : Va 

o

Gβ H =3.5×CIV_A+3.5×CCS_A-3.5×CIS_A Cu : Cu : Va : Va : Va 

o

Gβ H = 3.5×CCS_A+3.5×CIH_A-3.5×CIS_A +2

Cu : Cu : Va : Va : Va 

o

Gβ H = 3.5×CIS_A+TEGN −1

Cu : In : Se : Va : Va 

o

Gβ H = 3.5×VIS_A+TEGN −1

Va : In : Se : Va : Va 

o

Gβ H = 3.5×EIS_A+TEGN −1 −1

Va : In : Se : Va : Va 

o

Gβ H = 3.5×IIS_A+TEGN −1

In : In : Se : Va : Va 

o

Gβ H = 3.5×CCS_A+TEGN −1

Cu : Cu : Se : Va : Va 

o

Gβ H = 3.5×CIV_A+TEGN −1

Cu : In : Va : Va : Va 

o

Gβ H = 3.5×CIH_A+TEGN +2 −1

Cu : In : Va : Va : Va 

151
o

Gβ H = 3.5×CIS_A+TEGH +1

Cu : In : Se : Va : Va 

o

Gβ H = 3.5×VIS_A+TEGH +1

Va : In : Se : Va : Va 

o

Gβ H =3.5×EIS_A+TEGH −1 +1

Va : In : Se : Va : Va 

o

Gβ H = 3.5×IIS_A+TEGH +1

In : In : Se : Va : Va 

o

Gβ H = 3.5×CCS_A+TEGH +1

Cu : Cu : Se : Va : Va 

o

Gβ H = 3.5×CIV_A+TEGH +1

Cu : In : Va : Va : Va 

o

Gβ H = 3.5×CIH_A+TEGH +2 +1

Cu : In : Va : Va : Va 

o

Gβ H = 3.5×CIS_A+TEGH+TEGN −1 +1

Cu : In : Se : Va : Va 

γ-CIS o

Gγ = 4×CIS_A+44500 (Cu Cu : In : Se : Va : Va 1In5Se8) 

o

Gγ = 4×CCS_A Cu : Cu : Se : Va : Va 

o

Gγ = 4×VIS_A Va : In : Se : Va : Va 

o

Gγ = 4×IIS_A In : In : Se : Va : Va 

o

Gγ = 4×CIV_A Cu : In : Va : Va : Va 

o

Gγ = 4×EIS_A −1

Va : In : Se : Va : Va 

o

Gγ = 4×CIH_A +2

Cu : In : Va : Va : Va 

o

Gγ = 4×CCS_A+4×VIS_A-4×CIS_A Va : Cu : Se : Va : Va 

o

Gγ = 4×CCS_A+4×EIS_A-4×CIS_A −1

Va : Cu : Se : Va : Va 

o

Gγ = 4×CCS_A+4×IIS_A-4×CIS_A In : Cu : Se : Va : Va 

o

Gγ = 4×VIS_A+4×CIV_A-4×CIS_A Va : In : Va : Va : Va 

o

Gγ = 4×EIS_A+4×CIV_A-4×CIS_A −1

Va : In : Va : Va : Va 

o

Gγ = 4× IIS_A+4×CIV_A-4×CIS_A In : In : Va : Va : Va 

152
o

Gγ = 4×CIH_A+4×VIS_A-4×CIS_A +2

Va : In : Va : Va : Va 

o

Gγ = 4×CIH_A+4×EIS_A-4×CIS_A −1 +2

Va : In : Va : Va : Va 

o

Gγ = 4×CIH_A+4×IIS_A-4×CIS_A +2

In : In : Va : Va : Va 

o

Gγ = 4×CIV_A+4×CCS_A-4×CIS_A Cu : Cu : Va : Va : Va 

o

Gγ = 4×CCS_A+4×CIH_A-4×CIS_A +2

Cu : Cu : Va : Va : Va 

o

Gγ = 4×CIS_A+TEGN −1

Cu : In : Se : Va : Va 

o

Gγ = 4×VIS_A+TEGN −1

Va : In : Se : Va : Va 

o

Gγ = 4×EIS_A+TEGN −1 −1

Va : In : Se : Va : Va 

o

Gγ = 4×IIS_A+TEGN −1

In : In : Se : Va : Va 

o

Gγ = 4×CCS_A+TEGN −1

Cu : Cu : Se : Va : Va 

o

Gγ = 4×CIV_A+TEGN −1

Cu : In : Va : Va : Va 

o

Gγ = 4×CIH_A+TEGN +2 −1

Cu : In : Va : Va : Va 

o

Gγ = 4×CIS_A+TEGH +1

Cu : In : Se : Va : Va 

o

Gγ = 4×VIS_A+TEGH +1

Va : In : Se : Va : Va 

o

Gγ = 4×EIS_A+TEGH −1 +1

Va : In : Se : Va : Va 

o

Gγ = 4×IIS_A+TEGH +1

In : In : Se : Va : Va 

o

Gγ = 4×CCS_A+TEGH +1

Cu : Cu : Se : Va : Va 

o

Gγ = 4×CIV_A+TEGH +1

Cu : In : Va : Va : Va 

o

Gγ = 4×CIH_A+TEGH +2 +1

Cu : In : Va : Va : Va 

o

Gγ = 4×CIS_A+TEGH+TEGN −1 +1

Cu : In : Se : Va : Va 

153
Functions 

CIS_A= GCUFCC+2×GSE_S+GIN_S-226835 

VIS_A= GIN_S+2×GSE_S+20000 (298 ~ 6000 K) 

VIS_B= GIN_S+2×GSE_S-169043 

IIS_A= 2×GIN_S+GSE_S+94873 

CIV_A= GCUFCC+GIN_S+46154.3 

CCS_A= 2×GCUFCC+2×GSE_S-28502 

EIS_A= GIN_S+2×GSE_S-178812 

CIH_A = GCUFCC+GIN_S+43729 

DC_A= 2×GIN_S+2×GSE_S-2×GCUFCC-195049 

LCV = -27468 

LCE = -26218 

LCI = -134189 

LIC = -1039 

LSV = -28861 

LSH= -20000 

TEGN = 46474+233.51T-12.471TLN(T) 

TEGH = 46474+207.55T-12.471TLN(T) 

GCUFCC= -7770.45775+130.485222T-24.11239Tln(T) 

-.00265684T2

+1.29222833E-07T3+52477.8T-1 (298 ~1368 K) 

= -13309.7197+183.649837T-31.38TLN(T) (1358~3200) 

GIN_S= -6978.89011+92.3381153T-21.8386Tln(T) 

-.00572566T2

-2.12032167E-06T3-22906T-1 (298~430 K) 

 = -7032.8091+124.475108T-27.4562Tln(T)+5.4607E-04T2

 -8.36698333E-08T3

-211707.5T-1 (430~3800 K) 

154
GSE_S = -6657.653+92.539695T-19.14Tln(T) 

-.012295T2

+2.6766666E-06T3 (298~ 494 K) 

= -6657.653+92.539695T-19.14Tln(T)­

.012295T2

+2.6766666E-06T3 (494~760K) 

= -9059.16586+150.334216T-28.552Tln(T) (760~1500 K) 

Note : All the units are in J/mol 

155
8 Formation of α-CuInSe2 by Rapid Thermal Processing 

of a Stacked Binary Compound Bilayer 

8.1 Introduction 

Common methods of CIS absorber formation (e.g., physical vapor deposition (PVD) and 

selenization) require high-temperature growth conditions. The PVD needs precise control of the 

flux of each element. The temperature for the process is over 550 °C for the formation of CIS 

such as that used in the NREL three stage process. The selenization involves a number of 

reactions and the complex interdiffusion of intermediate phases affecting the controllability of 

the film quality. Care must be taken to prevent the formation of high melting temperature 

phases, such as Cu2-xSe and In2Se3, which will hinder the formation of CIS. Researchers use 

H2Se instead of selenium for selenization to solve the selenium dewetting problem and to avoid 

the formation of high melting temperature secondary phase. However, H2Se is highly toxic, and 

the selenization rate is slow. Thus the purpose of this work is to demonstrate a low temperature 

route for the formation of CIS by rapid thermal processing (RTP) of binary bilayer precursors. 

8.2 Background 

The ultimate goals of the research and development efforts are to create low cost, 

reproducible process with good process control. To accomplish these goals, fundamentally sound 

decisions must be made regarding the basic process and equipment to be employed, and then the 

entire process must be understood and characterized in terms of device behavior as a function of 

processing conditions. Among all the techniques, RTP is promising. Rapid implies high 

throughput as well as potentially lower thermal budget and safer process. The rapid heating 

feature can also give better control of the process kinetics. Most of all, a steeper temperature 

profile lowers substrate temperature. Lower thermal budget also give larger flexibility on the 

substrate selection. 

The Cu-In-Se phase diagram suggests that alternative precursor structures can be used for 

the selenization method for the CIS absorber formation. A scheme of the proposed precursor 

structure is shown in Figure 8-1. It consists of a slightly Se-rich In-Se or Cu-Se binary layer 

deposited at a low temperature on a Mo-glass substrate, followed by the deposition of a Se-rich 

Cu-Se or InSe binary layer, respectively. The stacked metal-Se bilayer precursor structure (CuSe/In-Se/Mo/glass) was rapidly deposited at 200 ̊C in a custom designed Migration Enhanced 

Epitaxy (MEE) system [8-1] that accommodated 8 2” by 2” substrates, and subsequently 

annealed using RTP. 

This structure was suggested by an examination of the Cu-In-Se phase diagram for 

conditions that show α-CuInSe2 in equilibrium with a liquid phase at low temperature. Precursor 

layer compositions were then selected on the likelihood of forming the liquid phase while 

avoiding high melting temperature compounds. The α-CuInSe2 phase is seen to lie on the tie line 

connecting the In2Se3 and Cu2Se phases, and it shows a solid solution range of several atomic 

156
percent [8-2]. A low-temperature single phase Se-rich liquid solution is located near the Se 

corner of the phase diagram. The triangle connecting the α-CuInSe2 phase boundary to the Se 

corner represents the α-CuInSe2 plus Se-rich liquid two-phase region and is the target for the 

overall film composition. Phase diagram calculations suggest that the Se-rich liquid phase can 

exist at a temperature as low as 211  ̊C [8-3]. They also suggest that the synthesized CIS will be 

Se-rich and excess Se can easily volatilize during annealing. Starting from binary metal-Se layers 

in contact promises to avoid certain reactions that yield high melting temperature compounds 

and to increase grain size through liquid phase assisted growth. The rapid thermal processing 

mitigates undesirable solid to solid transformations predicted at low temperatures. 

Why do we choose Cu:Se and In:Se layers with composition of ~ 1:1? The compounds 

CuSe and InSe have much lower melting temperature than Cu2Se and In2Se3 [8-3], and the liquid 

phase can be easily achieved by rapid thermal processing because the rapid heating can allow us 

to avoid the solid and solid transformation. So, the slightly Se-rich CuSe and InSe composition 

are the best pair for the formation of α-CIS. With the precursor structure of 

CuSe/InSe/Mo/Glass, the top CuSe can prevent the In and Se loss from the bottom InSe. The 

vice versa structure (InSe on top of CuSe) will provide more efficient use of liquid phase which 

is formed by CuSe phase and thus achieve large grain growth. 

8.3 Experimental apparatus 

The system used for absorber layer growth on binary compound-based precursor is a 

4100 HEATPULSE Rapid Thermal Processing system supplied by AG Associates [8-3]. The 

RTP system is located in a clean room at the Surge Area at the University of Florida. The 

system is a single-wafer, cassette-to-cassette rapid thermal processor, capable of processing in 

inert or corrosive ambient. Currently, only N2 is being used in the system. The HEATPULSE 

4100 system components include: Heating Chamber, Quartz Isolation Tube, System Controller, 

Robot, Gas Control Electronics, ULPA Filtration System, Cooling Water Manifold, and 

Temperature Measurement Instruments. The oven isolates the wafer to be processed on a quartz 

wafer tray inside a quartz tube located inside the actual heating chamber. The quartz isolation 

tube has a small opening in the rear for the introduction of gas used in wafer processing. This is 

the purge inlet fitting. Gas is exhausted through a hole in the door flange. The 208V, 1200W 

tungsten halogen lamps are arranged in a bank of 10 above and 11 below the isolation tube. The 

lamps are connected at each end to the triac plate mounted directly above the oven. The tungsten 

halogen lamps are controlled by triacs. The triacs are protected by an over-temperature sensor. 

The RTP system is controlled by software that is menu-driven. The computing and userinterface functions of the RTP system consist of a computer and a touch screen. The waferhandling system is comprised of the robot, robot controller, and connections for communications 

signals and vacuum. The robot provides automated loading and unloading of samples. Software 

modifications had been made to slow the speed of the robot movements. 

The gas box is mounted on the left side of the system, and it can be connected to the 

house exhaust system by the exhaust hood mounted on the top of the gas cabinet. The ULPA 

filter unit is located above the wafer handling area, and it provides a local clean environment for 

loading and unloading samples. The entire unit can be removed if filter replacement or other 

servicing is required. The cooling system contains oven chamber cooling, pyrometer cooling, 

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tube cooling, and electronic assembly cooling. The heating chamber cooling system is connected 

to house chilled water at 65°F. The pyrometer cooling is provided by the internal closed loop 

chiller at about 15°C. The wafer temperature is monitored by either an optical pyrometer or a 

type K thermocouple. The pyrometer or thermocouple sensing offers precise closed-loop 

temperature control. Currently, the wafer temperature is monitored by an optical pyrometer. 

The front view of the HEATPULSE 4100 RTP system is shown in Figure 8-2. 

The HEATPULSE 4100 RTP system offers a great deal of flexibility in designing 

time/temperature profiles. Variable ramp rate, time, and temperature for each step in a cycle, and 

multi-step cycles that incorporate more than one steady state period can be programmed in the 

RTP system. The ramp rate can be as high as 250°C/sec. The programmable range for the 

steady state process time is 1~600 sec. 

The short processing time and high ramp rate characteristics of this RTP can dramatically 

reduce volatilization losses during the reaction and provides liquid-phase assisted growth of the 

absorber layer. By carefully designing and choosing binary pairs for the precursor, the RTP can 

also help to avoid the high melting temperature phase formation. To understand the processing 

parameter effects on the reaction and crystallization kinetics, various temperature profiles were 

used by varying for example the temperature ramp rate, maximum temperature and soak time, 

and etc. Therefore, our objective is to utilize the processing advantages of RTP for the 

manufacture of CIS and CIGS solar cells while tailoring the precursor structure using the 

research PMEE reactor. 

8.4 Film Characterization 

The instruments used to study the film properties include XRD and SEM for structural 

characterization; EDS, AES, EPMA and ICP for composition analysis; I-V and QE for device 

performance measurements; and Deep Level Transient Spectroscopy for trap level and trap 

density study. 

8.4.1 Crystal Structure Determination by X-ray Diffraction (XRD) 

The XRD is the most widely used technique for individual crystalline phase identification 

and its relative amount determination. An atomic plane (hkl) is spaced dhkl apart in distance from 

one another. (hkl) are the Miller indices for an atomic plane. The XRD instrument used in this 

study is Philipe 3720. It uses X-ray wavelength of 1.54 Å Cu Kα source. When X-rays are 

scattered by the atomic plane in a crystal and the diffraction angle 2θhkl satisfies the Bragg’s law 

λ=2dhklsinθhkl, a diffraction peak can be observed. θhkl is the angle between the atomic planes 

and the incident (or diffracted) X-ray beam. The measured d-spacing in the diffraction pattern is 

compared with known standards in the JCPDS powder diffraction file to identify the present 

phases. If there are several phases in the films, the proportion of each phase can be determined 

from integrated intensities in the diffraction pattern. 

8.4.2 Topographical Information - Scanning Electron Microscopy (SEM) 

The SEM is a commonly used instrument for image analysis. A source electron beam is 

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focused into a fine probe and rastered over the surface of the sample. With appropriate detectors, 

secondary electron and backscattered images can be obtained to provide the surface 

topographical information. 

8.4.3 Local Composition Analysis - Energy Dispersive Spectroscopy (EDS) 

When an ionized atom returns to its ground sate, it can release a characteristic X-ray. 

The concentrations of all the elements in the sample can be given by the analysis of the 

histogram of counts versus X-ray energy. EDS is always attached to electron column 

instruments like the SEM. 

8.4.4 Composition Depth Profile - Auger Electron Spectroscope (AES) 

The electron bombardment on the surface will produce an inner shell vacancy for the 

atom. Once the inner shell vacancy is created in an atom, the atom may then return toward its 

ground state via emission of a characteristic X-ray or through a radiationless Auger transition. 

The number of Auger electrons from a particular element emitted from a volume of material is 

proportional to the number of atoms of the element in the volume. Therefore, the relative ratio 

of elemental concentration can be determined from the energy distribution output of the detector. 

The Auger spectroscope is combined with ion-beam sputtering to remove materials from the 

surface and to continue to monitor the composition and the chemistry of the remaining surface. 

By this means, the depth concentration profile is obtained. 

8.4.5 Local Composition Analysis - Electron Probe X-ray Microanalysis (EPMA) 

The physical basis of the EPMA is essentially the same as that of the EDS analysis. The 

only difference is the detector. The EDS uses the photoelectric absorption of the X-ray in a 

semiconductor crystal (silicon or germanium), with proportional conversion of the X-ray energy 

into charge through inelastic scattering of the photoelectron. The quantity of the charge is 

measured by a sophisticated electronic circuit linked with a computer-based multichannel 

analyzer to collect data. The EMPA detector is a focusing device and is based on the 

phenomenon of Bragg diffraction of X rays incident on a crystal. It has different crystals to 

detect a wide energy range. Therefore, The EPMA has higher peak-to-background ratio than 

EDS and can detect lower concentrations. 

8.4.6 Overall Composition Analysis - Inductively Coupled Plasma (ICP) 

In ICP, ions are generated in an inductively coupled plasma and subsequently analyzed in 

a mass spectrometer. The CIS film is dissolved by nitric acid prior to the analysis and is 

introduced to the plasma as a liquid. 

8.4.7 I-V Measurements 

A solar cell mainly consists of a diode, whose forward bias dark current-voltage 

relationship (i.e., I-V curve under no light illumination) can be expressed as 

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= 0 { ⎣ ( s ) II exp ⎡qV − IR / AkT ⎦

⎤ −1 } (8-1)

where I0 is the reverse saturation current and is proportional to absorber bandgap 

by exp (−E

g / nkT ), Rs is the sheet resistance, n is any integral number larger than 1 and A is the 

diode factor. 

Under illumination, 

I = I0 exp ⎡

⎣ ( − IR s )/ AkT ⎤⎦ { − } L qV 1 − I (8-2) 

where IL is the light generated current. The optimal bandgap for the output of maximum 

efficiency is 1.4~1.5 eV. Increasing bandgap Eg will decrease I0 and in turn increase the voltage 

obtainable from the cell. However, increasing Eg will also decrease the IL because less solar 

spectrum is absorbed. 

There are several important parameters other than light-generated current involved in a 

solar cell. Short-circuit current (Isc) is the output current of the solar cell when the external 

circuit is shorted, i.e., zero load resistance. Open-circuit voltage (Voc) is the output voltage of the 

solar cell when the external circuit is open, i.e., infinite load resistance. The Voc is also referred to 

as photovoltaic voltage. Im and Vm are the current and voltage for maximum power. The solar 

efficiency of the cell is defined as follows: 

η = IV / P (8-3) m m incident 

ImVm is related to Voc and Isc by the fill factor ff: 

ff = IV / sc IV (8-4) m m oc 

8.5 Experimental Results and Discussions 

8.5.1 Susceptor Design 

The design of the susceptor went through three stages due to problems encountered at 

each stage as shown in Figure 8-3. Susceptor I (Figure 8-3 (a)) is designed to solve the cracking 

problem when using a silicon wafer as the susceptor as evidenced by the optical microscopy of 

CIS absorber plan view shown in Figure 8-4. Unfortunately, In and Se losses occurred while 

using susceptor I according to Auger (Figure 8-7) and XRD (Figure 8-5) analysis. Thus, 

Suceptor II (Figure 8-3 (b)) was designed to solve this problem. However, non-uniform heating 

occurred with the increased sample size. Therefore, Susceptor III (Figure 8-3 (c)) has been 

designed to successfully solve this problem. 

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8.5.1.1 Susceptor I 

The SLG sample with sample size 0.5” by 0.5” was initially placed on the top of a piece 

of graphite wafer. After processing films with this design, the bottom of the sodium-lime glass 

was observed to have exceeded the glass transition temperature. The test samples were all 

cracked and adhered to the silicon wafer. Optical micrographs of the samples were taken before 

and after the RTP and are shown in Figures 8-4. Therefore, a 4-inch diameter graphite plate as 

shown in Figure 8-4 was substituted for the silicon wafer to increase the heat conductivity of the 

sample tray. Several RTP experiments were conducted using graphite wafer. The plane view 

optical micrograph shown in Figure 8-4 (c) shows that the cracking and sticking are avoided by 

using the graphite tray. 

Three samples with precursor structure CuSe/InSe/Mo/sodium-lime glass were 

processed. Samples were placed on a 0.5 mm thick piece of 4-inch diameter graphite with a N2 

purge flow originating from the back of the oven. The RTP annealing recipes for the 3 samples 

are summarized in Table 8-1. 

Table 8-1. RTP conditions for processing CuSe/InSe/Mo/SLG sample 

Sample Soak time Temperature Ramp rate 

1-1 70 s 550 °C 20 °C/s 

1-2 70 s 750 °C 20 °C/s 

1-3 140 s 550 °C 20 °C/s 

The processed films were characterized by the XRD, Auger, and optical microscopy. 

From the XRD data shown in Figure 8-5, it can be seen that after RTP, the α-CuSe phase 

disappeared and the α-CuInSe2 and Cu2-xSe phases formed. The relative intensity of the peak 

around 2θ = 26.560 decreased when the soak time increased from 70 s (sample 1-1) to 140 s 

(sample1-3). This is attributed to the longer soak time increasing the In and Se losses and thus 

increasing the extent of formation of the Cu2-xSe phase. It is likely that the In2Se species is 

responsible since it has a high vapor pressure. There is little difference in the XRD spectra when 

the temperature is increased from 550 °C (sample 1-1) to 750 °C (sample 1-2) for the same soak 

time. 

From the Auger data shown in Figure 8-7, In and Se loss is observed after RTP and as the 

temperature increases from 550 oC to 750 oC, there are more loss of these two elements. 

From the plane-view optical micrographs, as illustrated in Figures 8-6 (a)-(c), it can be 

seen that when the set temperature is at 750 °C (sample 1-2), the surface morphology of the 

processed film is much rougher than the one processed at 550 °C. 

The above preliminary experimental results indicated that a high RTP temperature should 

be avoided and the loss of In and Se needs to be reduced. Therefore, a new sample holder was 

designed using a piece of flat quartz as a cover. This design is termed susceptor II. 

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8.5.1.2 Susceptor II and Influence of RTP Parameters Effect Study 

A set of RTP experiments was performed on precursor samples (0.5” by 0.5”) with 

CuSe/InSe/Mo/Glass structure to determine the effects of the RTP processing variables 

(temperature, soak time, delay times, sample size, and preheat intensity) on the composition 

change and phase transition of the film using susceptor design II as shown in Figure 8-3 (b). In 

this design, a quartz lid is used to cover the sample to minimize volatilization losses. A typical 

RTP cycle is shown in Figure 8-8. 

Preheat Intensity Effect 

Two kinds of preheat intensity were employed, 0% and 25%, for the same process 

condition 600  ̊C, 60 s, 60  ̊C/s Delay I 60 s, and Delay II time 200 s. The XRD and ICP results 

are shown in Figure 8-9. There is no significant change in film properties when increasing the 

preheat intensity to 25% except a slight increase in peak intensity at 26.5°. This increase is due 

to the increase extent of formation of Cu2-xSe because even though Cu2-xSe peak overlaps with 

the CIS peak, Cu2-xSe usually has another peak at the shoulder of 112 peak position and the 

FWHM value for Cu2-xSe is less than that of CIS due to its higher crystallinity. The maximum 

preheat intensity for this RTP is 30%. Therefore, lower preheat intensity is preferred. 

Steady State Temperature 

For both soak times (Figure 8-10 and 11), when increasing the temperature there is more 

Se loss. The FWHM of the decreases of the 112 peak due to increase formation of Cu2-xSe. 

Therefore, a lower soak temperature is preferred. 

Ramp Rate 

When decreasing the ramp rate (Figure 8-12), the XRD pattern shows similar peak 

position, peak intensity and FWHM at 26.5°. There is a large decrease in Cu/In ratio when 

increasing the ramp rate while the Se composition remains the same. This means that a higher 

ramp rate will favor the formation of CIS and thus have more indium in the film. Therefore, a 

higher ramp rate is preferred. 

Soak Time 

At both a high temperature (Figure 8-13) and a low temperature (Figure 8-14), when 

decreasing the soak time, the FWHM decreases by 30%. This means that less CIS phase is 

formed in the film processed at a shorter soak time. There is no significant difference in 

composition when varying the soak time. Therefore, a longer soak time is preferred at both a 

high and a low temperature to form the CIS phase. 

In conclusion, the film cracking and elemental loss issues were largely resolved by using 

a graphite sample tray and quartz cover even though slight loss of In and Se is still observed. 

From the XRD patterns measured for the films before and after the RTP, it can be seen that both 

the α-CuInSe2 and Cu2-xSe phases are synthesized by the RTP from the binary bilayer precursors 

and the α-CuInSe2 phase exhibited a chalcopyrite structure as indicated by the presence of the 

162
101, 103, 211 and 105 reflections [8-4, 5]. The α-CuSe phase, which is the dominant phase in 

the precursor film, disappeared and reacted to form CuInSe2 or Cu2-xSe phases upon RTP. The 

narrow peaks indicate large crystal sizes. This suggests that the Cu-rich secondary phase is the 

competing phase in the binary bilayer process. However, it modifies the film morphology and 

results in large grain growth. From the XRD and ICP results, a higher ramp rate and a longer 

soak time have a significant positive influence on the α-CIS phase formation. A higher 

temperature will produce more loss in Se. Therefore, the starting RTP conditions for the 

following experiments on the precursor structure CuSe/InSe/Mo/Glass are: higher ramp rate, 

longer soak time, 0% preheat intensity, and lower steady temperature (around 500 °C). 

8.5.1.3 Susceptor III 

The RTP experiments were performed on new precursor series (growth run 230) with the 

structure CuSe/InSe/Mo/glass. For this sample set, the sample size was increased to 1” by 1” to 

allow cells to be fabricated. Experiments performed in the sample susceptor II (Figure 8-3 (b)) 

produced large lateral temperature gradients during RTP because of the relatively small size of 

the quartz cover. XRD spectra of the three different locations (see Figure 8-15) were taken, and 

the phase constitution of the three regions is shown in Figure 3-16. It can be seen that the Cu-Se 

secondary phase peak (2θ=21.3°) decreases in the center region as compared to the edge, 

consistent with a higher center temperature. The reaction in the center is thus more complete 

than at the edge. To improve heating uniformity, a new sample holder was designed (Figure 3-3 

(c)). In this design, a recess is milled in a piece of graphite stock that holds the sample. A 

second recess is milled at a larger diameter to accommodate a flat quartz cover. This design 

used for the RTP experiments was made to minimize Se loss and to achieve a uniform 

temperature. 

8.5.2 Phase Transition Study on CuSe/InSe/Mo/Glass Precursors 

A phase transition study was conducted on the precursor structure CuSe/InSe/Mo/Glass 

using susceptor III to find a suitable reaction pathway for the formation of α-CIS. 

Experimental 

The stacked metal-Se bilayer precursor structure (Cu-Se/In-Se/Mo/glass) was rapidly 

deposited at 200 ̊C in a custom-designed Migration Enhanced Epitaxy (MEE) system [8-1] that 

accommodated 8 2” by 2” substrates. A single precursor film was then annealed in a series of 4 

steps in the RTP system. Each anneal step exposed the sample to a different ambient by coating 

the glass cover with either InSe or Se. The RTP experiments were performed on 1” by 1” 

samples from the same growth run. 

The whole investigation consists of five steps as follows: 

Step I. Anneal using a clear quart cover 

Initial experiments were conducted using a simple quartz cover with the gas ambient 

established by volatilization of the sample. The RTP conditions of a set-point temperature 480 

°C, soak time 60 s, and ramp rate 60 °C/s were first tested. The XRD results, however, indicated 

163
incomplete formation of α-CuInSe2, with intermediate phases (predominantly InSe, In6Se7, Cu2­

xSe and CuSe2) remaining (see Fig. 8-17). Suspecting reaction limited processing, the set point 

temperature was raised to 600°C and the soak time increased to 120 s. The XRD results shown 

in Fig. 3-18, however, do not show a significant change in the annealed film constitution. The 

overall composition of the annealed sample (temperature 480°C, soak time 60 s, and ramp rate 

60 °C/s) was determined by the ICP analysis to be Cu 29.4%, In 23.6%, and Se 47.0 %. This 

composition is located in the three-phase region (liquid + Cu2-xSe + α-CuInSe2) according to the 

isothermal section of the phase diagram [8-2]. These results suggest that Se volatilization 

occurred during processing. To maintain the Se concentration during annealing, Se was added to 

the system by depositing a thin layer of either InSe or Se on the bottom of the quartz cover. 

Step II. Anneal using a quartz cover with an InSe film 

To modify the gas ambient, an1 μm film of InSe was deposited onto the quartz cover. 

The sample annealed in step I at RTP conditions 480 °C, 60 °C /s, and 60 s was again processed 

under the same conditions but with the InSe coated cover in place. As shown in the XRD pattern 

in Fig. 8-19, Cu2-xSe was still present. 

Step III. Anneal using a quartz cover with a Se film 

To increase the Se overpressure, a thin layer of elemental Se was next deposited on the 

cover and the same sample was once more processed under identical RTP conditions (480 °C, 60 

°C /s, 60 s). The XRD analysis showed an increase in the peak intensity at 2θ of 26.5 ̊ attributed 

to α-CuInSe2, suggesting that InSe reacts with Cu2-xSe in the presence of Se (see Fig. 8-20). In 

addition small amounts of Se-rich CuSe2 and In6Se7 were also apparent. 

Step IV. Anneal with a clear quartz cover 

To reduce or eliminate the Se-rich CuSe2 and In6Se7 phases, the same film was again 

annealed at the same conditions, but with no deposit on the underside of the cover plate. The 

results shown in Fig. 8-21 indicate that the final film is a single-phase α-CuInSe2. 

Discussions and Conclusions 

XRD analysis of the precursor film (see Fig. 8-22) reveals only α-CuSe in the starting 

Cu-Se layer and no apparent crystalline phase in the In-Se layer. After the step I annealing, 

crystallization of InSe is evident, while secondary Cu2-xSe is formed, presumably in the top layer 

as a result of Se loss. The compound α-CuSe will decompose to Cu2-xSe with loss of Se at 377 

 ̊C [8-5] according to: 

α − CuSe(s) → 1

2 β − Cu 2−xSe(s) + 14 Se 2 (8-5) 

The constitution did not change significantly when the same sample was again annealed 

using an InSe coated slide. Apparently, there was insufficient mass transfer between layers to 

allow measurable reaction according to: 

164
β−Cu2−xSe(s) +InSe(s) →α−CuInSe2 (s) (8-6) 

Although no significant reaction of β-Cu2-xSe(s) with InSe(s) occurred in steps I to III, 

the formation of α-CuInSe2 proceeded rapidly in step IV. When a Se ambient is present during 

anneal, the following overall reaction occurs: 

β − Cu 2−xSe(s) + 2InSe(s) + 1

2 Se 2 (g) → 2α − CuInSe 2 (s) (8-7) 

This reaction occurs rapidly with an activation energy of 25 kJ/mol as determined by a 

study of the selenization of Cu-In bilayers [8-6]. The apparent role of Se is to retain the CuSe 

phase by the reaction: 

β − Cu 2−xSe(s) +Se 2 (g) → CuSe(s) (8-8) 

The SEM micrographs (see Fig. 8-23) support the above observations and the proposed 

reaction mechanism. The micrograph of the precursor structure (see Fig. 8-25(a)) clearly 

delineates the Cu-Se/In-Se/Mo layered structure. After the first anneal (step I), the three-layers 

are still evident, but the microstructure has changed (see Fig. 8-25(b)). According to the EDS 

analysis, the layer next to Mo layer is InSe, and the surface layer is β-Cu2-xSe while α-CuInSe2 is 

formed between these two layers. Fig. 8-25(c) shows the final film structure after step IV. As 

desired, a single-phase α-CuInSe2 is now present with a grain size on the order of a few microns. 

Based on the above experimental observation, we conclude that Se control is an essential 

requirement for the RTP process of binary bilayer precursor, and a higher substrate temperature 

will be suggested to support the formation of a Se-rich liquid CuxSe phase in Cu-Se film to assist 

the growth. 

−Cu 2−x 

Se→ High T 

x β Se s () ⎯⎯ CuSe s () ⎯⎯⎯→Cu Se Liquid ( ) (8-9) 

The liquid CuxSe can help the diffusion of In atom and thus the following reaction will 

happen: 

Cu Se Liquid ( ) + InSe →α-CuInSe (8-10) x 2 

8.5.3 RTP Experiments on CuSe/InSe/Mo/Glass Precursors in PVD system 

As discussed in the previous section, a Se ambient is needed for the formation of CIS, 

and a higher temperature is suggested to help the diffusion of In assisted by the formation of 

liquid CuxSe phase. Therefore, the following three runs were performed under Se over pressure

of 5 Å/s control by RTP on CuSe/InSe/Mo/Glass binary bilayers in the PVD system at NREL: 

Run 1. 500 °C, 2.5 minutes 

Run 2. 550 °C, 4.5 minutes 

Run 3. 550 °C, 10 minutes 

165
A safety limit on the PVD system restricted the ramp rate to 3 °C/s for all three 

experiments. This ramp rate is about 10 times faster than the regular ramp rate used in the 

normal PVD process. As discussed in the background section, a faster ramp rate enhances the 

reaction of these two layers by forming a liquid phase and results in a lower substrate 

temperature and shorter processing time. A conventional RTA system can have a ramp rate up to 

120 oC/s. Therefore, a conventional RTA system should allow the formation of CIS at lower 

process temperature and shorter processing time compared to the PVD system. 

As seen from the XRD pattern in Figure 8-22, only the CuSe phase is present in the 

precursor (before RTP). The bottom InSe phase does not show any peaks because it is in the 

amorphous state. The SEM cross section of the precursor (Figure 8-24b) shows the separate 

InSe and CuSe layer of thickness about 1 μm each. The plane view of the precursor (Figure 8­

24a) shows the hexagonal CuSe grains. The shoulder near (112) peak in the XRD pattern 

(Figure 8-25) of the film after Run 1 indicates the formation of the high temperature Cu2-xSe 

phase. The bright dots in the plane view micrograph indicated by an arrow in the Figure 8-26a 

points out the Cu2-xSe phase. The cross-sectional micrograph (Figure 8-26b) shows the small 

grain site in the bottom portion of the layer because of incomplete nucleation of CIS. It is 

believed that these small grains are not well nucleated CIS grains and that the bottom InSe is 

completely consumed as evidenced by absence of InSe secondary phase by XRD (usually InSe 

will crystallize upon heating). This also suggests that In diffusion in CuSe is slower than Cu 

diffusion into InSe. This is because the In atom is larger than the Cu atom, and InSe is 

amorphous which makes the Cu diffusion much easier. 

EPMA is a useful technique to determine the existence of any Cu rich secondary phases 

because it probes the upper layer composition, and Cu-rich secondary phases tend to segregate 

near the surface [8-6]. The EPMA composition for this film after Run 1 is Cu: 27.9 at.%, In 24.1 

at.%, Se 48.0 at.% yielding a Cu/In ratio of 1.16. This result is consistent with the existence of a 

Cu-rich secondary phase. Therefore, the temperature was increased to 550 °C and the soak time 

increased to 4.5 min in Run 2 to enhance In diffusion and thus formation of CIS in the top CuSe 

layer. The XRD pattern (Figure 8-27) shows the single CIS phase and the disappearance of Cu2­

xSe on the surface as suggested by the SEM plane view micrograph (disappearance of small 

surface grains in Figure 8-28(a)). The grain size in the bottom layer enlarged (Figure 8-28(b)). 

The composition as determined by EPMA for film after run 2 is Cu 23.5 at.%, In 26.5 at.% and 

Se 50.0 at.% with Cu/In ratio of 0.887 and this ratio lies in single phase CIS region. This 

suggests that the top layer is also fully converted to the CIS phase. To promote CIS grain growth 

in the bottom layer, a longer soak time, 10 min was tried (Run 3). From the SEM picture crosssection micrograph shown in Figure 8-29(b), a relatively homogenous grain size is evidence 

throughout the film, and the plane view SEM micrograph shows highly textured CIS grains as in 

Figure 8-29(a) and this is the highly oriented CIS 112 peak. 

I-V Measurements 

The next step was to fabricate cells from the absorbers synthesized in Runs 1-3. The I-V 

measurement for the first device is shown in Figure 8-30 (without Se overpressure after Run 1). 

The device is shorted apparently due to the presence of the conducting Cu2-xSe phase. For Run 

2 with XRD and SEM suggesting single phase CIS with two different grain size layers, the 

efficiency is 2.78 %. Further increase of grain size (Run 3) produced an increased efficiency of 

166
5.08 %. The Jsc for this 5 % cell is 34.65 mA/cm2

, which is a good value for CIS. Fill factor and 

Voc, however, are very low. A high sheet resistance contributes to the low fill factor. This can be 

improved by adjusting the Cu:In:Se composition ratio and thus the doping level, especially Cu to 

In ratio and by creating a composition gradient to achieve the bandgap grading with the 

incorporation of Ga. The Voc is largely affected by the mid gap defects in the space charge 

region. To detect the defect level in this 5 % cell, a Deep Level Transient Spectroscopy (DLTS) 

measurements were performed on this 5 % cell. 

8.5.4 RTP Experiments on InSe/CuSe/Mo/Glass Precursors 

A preliminary study of samples with the In-Se binary layer on the top of the Cu-Se layer 

(precursor structure InSe/CuSe/Mo/glass) was undertaken. Samples with this structure, when 

processed at 450 to 600 °C, showed a rough surface morphology after the RTP as shown in 

Figure 8-31 (a). This is believed to be a result of poor adhesion of the bottom layer to Mo layer. 

Therefore, a seed thin layer of InSe was first deposited on the Mo/Glass substrate to give 

structure InSe (1 μm)/CuSe (1 μm)/InSe(500 Å)/Mo/Glass. This structure greatly improves the 

surface morphology as seen in the optical micrograph as shown in Figure 8-31(b). The structure 

and process conditions however need to be optimized to make device quality films. 

8.6 Conclusions 

An alternative low temperature route to CuInSe2 formation was demonstrated using RTP 

on the stacked precursor structure Cu-Se/In-Se/Mo/glass. The influence of various ambient 

compositions on phase transformations after RTP was determined. A short time annealing of 60 s 

at 480  ̊C in a Se ambient generated by volatilization of a thin Se layer on the sample cover, 

followed by a second anneal using a clear cover, was able to produce single-phase α-CuInSe2 

with relatively large grain size. The high melting temperature secondary phase Cu2-xSe(s) does 

not react with InSe without Se overpressure. Based on these results, a 5.08% device with Voc = 

0.296 V, Jsc = 34.65 mA/cm2

 and fill factor = 49.54% was fabricated using the absorber that is 

obtained by RTP on the CuSe/InSe/Mo/Glass precursor under Se control in a PVD system. It is 

worth noting that this 5.08% cell was achieved without full optimization of the process. 

Therefore, the RTP on binary bilayer process is promising. 

8.7 References cited 

[8-1] B.J. Stanbery, C.-H. Chang, S. Kim, S. Kincal, G. Lippold, S.P. Ahrenkiel, L. L. Kerr, 

T.J. Anderson, M.M. Al-Jassim, Self Organized Processes in Semiconductor Alloys, 

MRS Symposium Proceedings, Vol. 583, 195-200, 2000. 

[8-2] T. Gödecke, T. Haalboom, and F. Ernst , Zeitschrift für Metallkunde, Vol. 91, No. 8 

pp.622-662, 2000. 

[8-3] C-H Chang, Processing and Characterization of Copper Indium Selenide for 

Photovoltaic Applications, Ph.D. Dissertation, University of Florida, 2000. 

167
[8-4] R. Friedfeld, R.P. Raffaelle, and J.G. Mantovani, Solar Energy Materials and Solar 

Cells, Vol. 58, pp. 375-385, 1999. 

[8-5] I. Forbes, F.O. Adurodija and M.J. Carter, Proceedings of the 6th International PVSEC, 

New Delhi, pp.1965-1968, 1996. 

[8-6] N. Orbey, H. Hichri, R.W. Birkmire and TWF Russell, 25th IEEE Photovoltaic Specialist 

Conference, May 13-17, Washington, D.C. pp. 981-982, 1996 

CuSe (at 50-55% Se) 

InSe (at 50-55% Se) 

RTP 

or α-CuInSe2 

InSe (at 50-55% Se) 

CuSe (at 50-55% Se) 

Figure 8-1. Novel precursor structure consisting of stacked metal-Se layers. 

168
(a) 

(b) 

Precursor film 

N2 

Sample susceptor 

Quartz tray 

Quartz tube 

Heating lamp 

Figure 8-2. AG Associates Heatpulse 4100 RTP system (a) front view, (b) 

schematic of the oven. 

169
(a) 

(b) 

(c) 

Figure 8-3. Susceptors design schematic (a) I, (b)II, and (c) III. 

170
(a) Trial sample before RTP. (b) Trial sample after RTP by (c) Trial sample after RTP by 

using silicon wafer as the using graphite wafer as the 

susceptor. susceptor. 

Figure 8-4. Comparison of film optical micrograph when using Si or graphite as 

the susceptor at magnification 10x.. 

8000 Mo 

CIS (112)/Cu2-xSe 

7000 CIS (220)/Cu2-xSe

CIS (312)/Cu2-xSe 

6000 CIS (103) After RTP 550 oC, 140 s 

α-CuSe (sample1-3) 

5000 

4000 After RTP 750 oC, 70 s 

(sample1-2) 

3000 

2000 After RTP 550 oC, 70 s 

(sample1-1) 

1000 

0 

Before RTP 

10 20 30 40 50 60 70 

2-Theta 

Figure 8-5. XRD pattern comparison for the films at different RTP conditions 

(Susceptor I).. 

Counts 

171
(a) Plane-view optical micro- (b) Plane-view optical micro- (c) Plane-view optical micrograph of sample 1-0 before graph picture of sample 1-1 graph picture of sample 1-2 

RTP 

Figure 8-6. Optical micrographs of 1-0, 1-1 and 1-3 samples at magnification 

10x.. 

172
0.1 

0.2 

0.5 

0.6 

0.7 

Se 

Cu 

In 

Figure 8-7 (a). Auger depth profile of sample 1-0 before RTP [8-3]. 

1-1 No Moly 1-3 No Moly 

0.8 

0.7 

0.6 Se 

Atomic Concentration (%)

Atomic Concentration (%) 0.5 O 

C 0.4 

O 

C 

Cu 

0.3 In

Se

Cu Cu 

In 0.4 

Se 

0.3 

0.2 

In 0.1 

0 

0 5 10 15 20 25 30 35 40 45 0 

Sputter Time (min) 0 5 10 15 20 25 30 35 40 45 

Sputter Time (min) 

Figure 8-7 (b). Auger depth profile of sample Figure 8-7 (c). Auger depth profile of sample 

1-1. 1-3. 

173
Figure 8-8. Typical Cycle time for Rapid Thermal Processing. 

174
Counts 

6000 

Mo 

α-CIS(112)/Cu2-xSe 

FWHM for peak at 26.5° =0.14 

5000 

α-CIS(220/204) 

ICP: Cu: 23.7%, In: 25.7% , Cu2-xSe 

4000 Se 50.6% α-CIS(116/312) 

α-CIS(101) α-CIS(211) Preheat intenisty 25% 

α-CIS(103)

3000 

2000 

FWHM for peak at 26.5° =0.14 

1000 ICP: Cu: 23.9%, In: 25.8% , 

Se 50.3% 

Cu2-xSe 

Preheat Intensity 0% 

0 

0 10 20 30 40 50 

2 Theta 

60 70 80 

Figure 8-9. The influence of preheat intensity on the phase formation (XRD 

pattern) and film composition (ICP). 

175
Counts 

16000 

α-CIS(112)/Cu2-xSe FHMW for peak at 26.5° = 0.14 

Mo 15000 

ICP: Cu: 24.2%, In: 26.0% , Cu2-xSe 

 Se 49.7% α-CIS(220,204) 

/Cu2-xSe 14000 α-CIS(116/312) α-CIS(211) /Cu2-xSe 

α-CIS(101) α-CIS(103) 600 °C 

13000 

12000 

ICP: Cu: 23.9%, In: 25.4% , 

Se 50.7% FHMW for peak at 26.5° = 0.16 11000 

10000 

480 °C 

9000 

8000 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-10. The effect of soak temperature at shorter soak time (XRD pattern 

and film composition by ICP). 

176
Counts 

1.4E+04 

α-CIS(112)/Cu2-xSe 

1.3E+04 ICP: Cu: 24.4%, In: 25.2% , Mo FHMW for peak at 26.5°= 0.24 

Se 50.4% Cu2-xSe 

α-CIS(220,204) 

/Cu2-xSe 1.2E+04 α-CIS(116/312) α-CIS(211) /Cu2-xSe 480 °C α-CIS(103) 

1.1E+04 

1.0E+04 

9.0E+03 ICP: Cu: 24.5%, In: 25.9% , 

Se 49.6% FHMW for peak at 26.5° = 0.2 

8.0E+03 

600 °C 

7.0E+03 

6.0E+03 

0 10 20 30 40 50 60 70 

2 Theta 

Figure 8-11. The effect of soak temperature at longer soak time (XRD pattern 

and film composition by ICP). 

177

80 
Counts 

9000 

α-CIS(112)/Cu2-xSe 

Mo 8000 

ICP: Cu: 24.7%, In: 24.9% , FWHM for peak at 26.5° =0.14 

Se 50.4% 

α-CIS(220,204)/Cu2-xSe 

Cu2-xSe 

α-CIS(116/312) 

7000 

6000 /Cu2-xSe 

Ramp Rate 20°C/s α-CIS(101) α-CIS(103) α-CIS(211) 

5000 

4000 

3000 

FWHM for peak at 26.5° =0.14 

ICP: Cu: 23.9%, In: 25.8% , 

Se 50.3% 

2000 

1000 Ramp Rate 60°C/s 

0 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-12. The effect of ramp rate (XRD pattern and film composition by 

ICP). 

178
Counts 

16000 

α-CIS(112)/Cu2-xSe Mo FHMW for peak at 26.5° = 0.14 

15000 ICP: Cu: 23.9%, In: 25.4% , Cu2-xSe

Se 50.7% α-CIS(220,204) 

/Cu2-xSe 

14000 α-CIS(116/312) α-CIS(211) 

α-CIS(103) /Cu2-xSe 60 s α-CIS(101) 

13000 

12000 

11000 

10000 

ICP: Cu: 24.4%, In: 25.2% , FHMW for peak at 26.5° = 0.2 9000 Se 50.4% 

8000 

120 s 

7000 

6000

0 10 20 30 40 50 60 70

2 Theta 

Figure 8-13. The influence of soak time at higher steady state temperature on 

phase formation (XRD pattern) and film composition (ICP). 

179

80 
Counts 

14000 

α-CIS(112)/Cu2-xSe 

13000 

Mo FHMW for peak at 26.5°= 0.24

ICP: Cu: 24.5%, In: 25.9% , 

Se 49.6% Cu2-xSe

α-CIS(220,204) 

/Cu2-xSe 12000 

α-CIS(116/312) α-CIS(211) α-CIS(103) /Cu2-xSe 

120 s α-CIS(101) 

11000 

ICP: Cu: 24.2%, In: 26.0% , 

Se 49.7% FHMW for peak at 26.5° = 0.16

10000 

60 s 

9000 

8000 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-14. The influence of soak time at lower steady state temperature on

phase formation (XRD pattern) and film composition (ICP). 

180
2 

3 

1 

Figure 8-15. Non-uniformly heated surface. 

181
1. InSe(100) Comparisonof three regions for non-uniformly heated surface 2. CIS(112) 

3. CIS(220)/(204) 

7 4. CIS(116)/(312) 12000 

5. Cu2Se(111) 

2,5 6. Cu2Se(220) 10000 

7. Cu2Se(311) 

9 3,6 8. CuSe(006) 8000 

1 4,7 9. Mo (110) Region 3 

6000 

4000 

2,5 

9 3,6 

1 4,7 Region 2 

2000 9 

1 2,5 3,6 4,7 Region 1

9 Precursor

Counts 

0 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-16. XRD pattern of different regions of non-uniform heated surface. 

182
(a) 

β-Cu2Se 

α-CuInSe2+β-Cu2Se 

InSe 

(b) 

0 

500 

1000 

1500 

2000 

2500 

Counts 

InSe 

α-CIS/Cu2SeH 

α-CIS/Cu2SeH 

Mo 

Cu2SeH 

In Cu2SeH 6Se7 

Cu2Se 

α-CIS/Cu2SeH 

Cu2Se 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-17. Film after step I: 480  ̊C, 60 s, 60  ̊C/s (a) structure (b) XRD 

pattern. 

183
(a) 

β-Cu2Se 

α-CuInSe2+β-Cu2Se 

InSe 

(b) 

0 

500 

1000 

1500 

2000 

2500 

Counts 

InSe 

α-CIS/Cu2SeH 

α-CIS/Cu2SeH 

Mo 

Cu2SeH 

In Cu2SeH 6Se7 

Cu2Se 

α-CIS/Cu2SeH 

Cu2Se 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-18. Film using higher temperature for step I: 550  ̊C, 120 s, 60  ̊C/s 

(a) structure (b) XRD pattern. 

184
(a) 

α-CuInSe2+InSe+β-Cu2Se 

InSe 

InSe Vapor 

(b) 

0 

500 

1000 

1500 

2000 

2500 

Counts 

α-CIS/Cu2SeH 

α-CIS/Cu2SeH 

InSe 

Mo 

α-CIS/Cu2SeH 

Cu2Se α-CIS 

Cu2Se 

Cu2Se 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-19. Film after step II: 480  ̊C 60 s, 60  ̊C/s with InSe lid (a) structure 

(b) XRD pattern. 

185
(a) 

αCuInSe2+Cu2Se+ 

In6Se7+CuSe2 

Se Vapor, Low T 

(b) 

0 10 20 30 40 50 60 70 80 

2 Theta 

Figure 8-20. Film after step III: 480  ̊C, 60 s, 60  ̊C/s with Se lid (a) proposed 

structure (b) XRD pattern. 

0 

500 

1000 

1500 

2000 

2500 

3000 

3500 

4000 

Counts 

InSe↓

α-CIS?(112)/Cu2SeH) 

α-CIS?(204)/Cu2SeH 

Mo 

α-CIS?(312) 

α-CIS↑

In6Se7 

Cu2Se/CuSe2 Cu2SeH 

CuSe2 

Cu2Se 

186
(a) 

α-CuInSe2 + very small amount 

InSe 

(b) 

0 10 20 30 40 50 60 70 

2 Theta 

Figure 8-21. Film after step IV: 550  ̊C, 70 s, 60  ̊C/s with clear lid (a) structure 

(b) XRD pattern. 

0 

1000 

2000 

3000 

4000 

5000 

6000 

7000 

Counts 

α-CIS (112) 

Mo 

α-CIS (204) 

α-CIS (312) 

α-CIS α-CIS α-CIS 

α-CIS InSe α-CIS 

187

80 
(a) 

InSe 

CuSe 

(b) 

0 10 20 30 40 50 60 70 

2 Theta 

Figure 8-22. Precursor (a) structure (b) XRD pattern and composition data. 

0 

5000 

10000 

15000 

20000 

25000 

30000 

35000 

40000 

45000 

50000 

Counts 

CuSe 

Mo 

CIS CuSe 

CuSe InSe 

CuSe 

Composition from ICP: 

Cu 24.28% In 26.43% Se 49.28% 

Cu/In ~ 0.92 

188

80 
(a) 

(b) 

(c) 

Figure 8-23. SEM cross-sectional micrograph of (a) precursor, (b) after step I, 

and (c) after step IV. 

Glass 

Mo 

InSe 

CIS 

CuSe 

Mo 

Glass 

InSe 

CuSe 

Mo 

Glass CuInSe2

189
(a) 

(b) 

Glass 

Mo 

InSe 

CuSe 

Figure 8-24. SEM micrograph of the precursor (a) plane view (b) cross-section. 

190
Counts 

18000 

CIS (112) 

16000 

14000 

12000 

10000 

8000 

6000 

Cu2-xSe(111) 

4000 

Mo CIS(204/220) 

2000 CIS(116) 

CIS(100) CIS(211) CIS 332/413 CIS(103)

0 

0 10 20 30 40 50 60 70 80 90 2 Τheta 

Figure 8-25. XRD pattern after RTP at 500 oC for 2.5 min (Run 1). 

191
(a) 

Cu2-xSe 

(b) 

Cu2-xSe 

Small Grain CuInSe2 

Larger Grain CuInSe2 

Figure 8-26. SEM micrograph of film after RTP at 500 oC for 2.5 min (Run 1) 

(a) plane view (b) cross-section. 

192
10000 

Counts 

12000 

CIS (112) 

8000 

6000 

4000 

MoCIS(204/220) 

2000 CIS(116) CIS(100) CIS(211) CIS 332/413 

CIS(103) 

0 

0 10 20 30 40 50 60 70 80 90 

2 Theta 

Figure 8-27. XRD pattern after RTP at 550 oC for 4.5 min (Run 2).. 

193
(a) 

CIS 

(b) 

Developed CIS 

CIS in development 

Figure 8-28. SEM micrograph of film after RTP at 550 oC for 4.5 min (Run 2) 

(a) plane view (b) cross-section. 

194
(a) 

(b) 

Figure 8-29. SEM micrograph of film after RTP at 550 oC for 10 min (Run 3) 

(a) plane view (b) cross-section. 

195
I-V Curve Voc 

(V) 

Jsc 

(mA/cm2) 

Isc 

(mA) 

FF 

(%) 

Eff 

(%) 

100

80

60

40

20

0

-20

-40

current dd ensity [ mA/cm^2]

-0.4

100

80

60

40

20

0

- 2 0

- 4 0

current d ensi ty [ m A/cm^2]

-0.4

100

80

60

40

20

0

- 2 0

- 4 0

current d ensi ty [ m A/cm^2]

-0.4

0.0

Voltage [ V]

 F0300/As dep without Se

0.0

Volt age [V ]

 F0303// 500 2.5 min /As dep

0.0

Voltage [V ]

 F0301/As dep 550°C/ 4.5 mins

0.4

0.4

0.4

0.8

0.8

0.8

0.025 

0.024 

0.262 

-0.48 

-4.34 

-30.36 

-0.21 

-1.86 

-13.05 

23.42 

24.25 

35.09 

0.000 

0.025 

2.79 

1 0 0

80

60

40

20

0

-20

-40

current dd ensi ty[ mA/cm^2]

- 0 .4 0 . 0

V o ltag e [V]

 M 00 0 0- 40 2 /a s d ep 55 0 °C/ 10 m in s

0 . 4 0 . 8

0.296 -34.69 -14.9 49.54 5.08 

Figure 8-30. Figure 3-30. I-V curves for the cells fabricated after without Se control and 

Runs 1-3 with Se control 

196 
(a) 

(b) 

Figure 8-31. Optical plane view micrograph absorber film after RTP process on 

inverted structure precursor (a) InSe/CuSe/Mo/Glass (b) 

InSe/CuSe/InSe/Mo/Glass at magnification 10x structure (b) XRD pattern. 

197
9 Investigation of Pulsed Non-Melt Laser Annealing on 

the Film Properties and Performance of Cu(In,Ga)Se2 

Solar Cells. 

9.1 Accomplishments 

CIGS cells contain a thin (~ 50 nm) n+ CdS buffer layer deposited on a thick (~ 1 to 2 

μm) p-type CIGS absorber. Although the exact location of the electrical junction relative to the 

metallurgical one is not known, it does lie near the surface of the absorber. Furthermore, 

analyses of CIGS film characterization results and cell performance data suggest that interfacial 

recombination near the CdS/CIGS metallurgical junction is a major limitation to the device 

performance. The possibility of repairing damage near this shallow junction, while preserving 

the composition gradients in the bulk CIGS films motivated exploration of NLA treatment. In 

this work CIGS absorber films were annealed using a 248 nm pulsed laser beam and the 

properties of the films and subsequent device performances were compared to untreated films. 

The CIGS samples used in this study were provided by the National Renewable Energy 

Lab (NREL) and Shell Solar Inc. (SSI), and both sample sets were grown on the Mo- coated 

soda-lime glass substrates. For the Hall-effect measurements, however, the CIGS films were 

grown on the insulating soda-lime glass (SLG) substrates. A thin (50 nm) CdS buffer layer was 

deposited on the CIGS samples using Chemical Bath Deposition (CBD) at 75 o

C for 30 min. The 

NLA treatment was carried out using a pulsed 248 nm line derived from a KrF excimer laser 

system. In this study the laser pulse width was fixed at 25 ns and the energy density was varied 

in the range 20 to 60 mJ/cm2

 and the number of pulse cycles in the range 5 to 20. To study the 

effect of NLA on the CIGS film properties, NLA treatments were performed on the CIGS films 

with and without CdS buffer layers. After NLA treatments, a ZnO window layer was deposited 

on the CdS/CIGS/Mo/SLG sample by RF sputtering. Subsequent metallization (Ni-Al front 

contact grids) was carried out by e-beam evaporation through a shadow-mask. Finally, finished 

devices were produced by cutting the sample into separate cells with 0.429 cm2

 active area and 

attaching wires with In bumps on the Mo-coated glass substrates for back contacts. The 

performance of these cells was then tested. 

The non-destructive Dual Beam Optical Modulation (DBOM) characterization method 

was used to measure the effective carrier lifetimes in the CIGS absorbers [9-4, 9-5], and to 

evaluate the effect of NLA treatment on the performance of CIGS cells. Within the sensitivity of 

the DBOM technique, the effective carrier lifetime was found to increase for CIGS films on 

Mo/SLG annealed using 5 cycles of laser pulse with an energy density in the range 30 to 60 

mJ/cm2

. The results summarized in Table 9-1 show that a low energy NLA treatment can 

increase the effective carrier lifetime of the annealed samples and improve the performance of 

CIGS cells. 

XRD patterns of CIGS and CdS/CIGS samples before and after NLA treatments show 

substantial narrowing of the diffraction peaks that belong to CIGS, as shown for two samples in 

Figure 9-1. These results can be interpreted as an improvement of the layer’s crystallinity after 

198
NLA treatment, which is consistent with the observed increase of the grain size (SEM) and the 

effective carrier lifetime in the CIGS films. 

Table 9-1 Effective minority carrier lifetime (τ) of CIGS films deposited on Mo 

coated soda-lime glass substrates before and after NLA treatment as measured by 

DBOM [9-4,9-5]. 

Sample # τ before NLA (ns) τ after NLA (ns) NLA condition 

CIGS-S1 1.77 4.87 30 mJ/cm2

, 5 pulses 

CIGS-S2 2.82 3.39 40, 5 

CIGS-S3 4.11 5.43 50, 5 

CIGS-S4 4.51 6.31 60, 5 

Figure 9-1. XRD spectra before and after NLA treatments for (a) a CIGS and (b) a 

CdS /CIGS sample. 

199
Figure 9-2. Surface morphology of CIGS films (a) without and (b) with NLA 

treatments at an energy density of 55 mJ/cm2 (SEM images with magnification of 

6000x). 

As illustrated by SEM micrographs shown in Figure 9-2, the surface morphology and 

apparent grain size increased upon laser annealing. This result suggests that the energy density 

used was sufficient to cause atomic rearrangement in the near surface region, and thus the 

potential exists to modify the atomic defects in the near surface region. 

Based on the encouraging DBOM results from the initial NLA treated CIGS samples, a 

second set of experiments were performed in which the energy density and number of the 

incident laser pulse were varied. Hall-effect measurements were made on the CIGS samples 

without a Mo layer prior and after annealing to determine the effects of NLA treatment on the 

carrier concentration, carrier mobility, and sheet resistivity of CIGS films. 

Four CIGS films deposited on the glass substrate underwent pulsed NLA treatments at 

room temperature. The annealing conditions and the results of Hall-effect measurements (using 

an MMR Hall and Van Der Pauw measurement system) before and after the NLA treatments are 

summarized in Table 9-2. The results show a significant increase in the values of Hall mobility 

and a decrease in film resistivity following NLA treatment. The carrier mobility in the NLA 

treated CIGS films were 3 to 4 times greater than the values before annealing. Although the hole 

concentration decreased slightly after annealing, film resistivities decreased by 72 and 64% for 

CIGS films (samples H1 and H2) treated at an energy density of 20 mJ/cm2

 with 10 and 20 laser 

pulses, respectively. At an energy density of 40 mJ/cm2

, the film resistivity was reduced by more 

than 95% from that of the non-annealed samples (H3 and H4). Thus, both the energy density and 

the number of laser pulse play an important role in determining the resistivity of NLA treated 

CIGS absorber layers. 

200
1 

2 

3 

4 

5 

6 

Table 9-2 Results of Hall-effect measurements for CIGS films deposited on the 

glass substrates before and after NLA treatments. 

Sample # Hole-concentration* 

(1016/cm-3) 

Hall-mobility* 

(cm2/V-s) 

Resistivity* 

(Ω-cm) NLA condition 

CIGS-H1 0.53/0.445 8.89/37.6 133/37.3 20 mJ/cm2

, 10 pulses 

CIGS-H2 2.9/2.43 0.93/2.98 235/86.4 20, 20 

CIGS-H3 4.3/1.8 1.54/6.1 94/2.67 40, 10 

CIGS-H4 7.1/3.3 0.60/2.8 148/4.64 40, 20 

* Data taken before/after NLA treatments. 

Six CdS/CIGS samples treated by NLA with selected energy density and pulse number 

plus one control sample without NLA treatment were fabricated into finished devices. There 

were five cells on each device and each cell was tested using a photo- J-V system at room 

temperature. The NLA conditions and averaged photo- J-V results of these devices are 

summarized in Table 9-3. Note that the data in Table 9-3 were averaged from the measured 

values of five cells on each device and the standard deviation computed on the 5- cell 

measurements. In these experiments, improvements were found in the photo- J-V results of the 

annealed devices under 30mJ/cm2

, using 5 or 10 pulses (CIGS-D1 and CIGS-D2), as compared 

to results from the control sample (Ctrl-D). The data also show degradation in cell performance 

parameters such as open circuit voltage (Voc), short-circuit current (Jsc), Fill Factor (F.F.), and 

cell efficiency (ηc) of the devices laser annealed with energy densities higher than 30 mJ/cm2

. It 

was also found that for devices annealed at the same energy density, the fill factor and 

conversion efficiency of the 5-cycle annealed devices are higher than those of 10-cycle annealed 

devices. These results suggest that an optimal NLA energy density and pulse number are 

approximately 30 mJ/cm2

 and 5 pulses, respectively. 

Table 9-3 Cell performance parameters (Voc, Jsc, F.F. and ηc) from the measured 

photo-J-V characteristics of NLA-treated CIGS solar cells 

Device # Voc (V) Jsc (mA/cm2) F.F. (%) ηc (%) NLA conditions 

CIGS-D 0.576 31.60 66.50 12.07 30 mJ/cm2

, 5 pulses 

CIGS-D 0.568 31.25 63.11 11.12 30, 10 

CIGS-D 0.497 26.79 55.29 7.368 40, 5 

CIGS-D 0.493 26.71 53.08 7.040 40, 10 

CIGS-D 0.436 25.56 51.80 5.773 50, 5 

CIGS-D 0.433 25.23 47.77 5.201 50, 10 

Ctrl-D 0.543 30.05 55.74 9.064 None 

201
To better understand the specific differences between the control cell and the two 30 

mJ/cm2

 NLA treated cells, dark J-V and quantum efficiency (Q-E) measurements as well as 

theoretical calculations of device parameters were made for selected cells from each of the three 

devices (Ctrl-D, CIGS-D1 and CIGS-D2). 

The J-V characteristics for a p-n junction solar cell under uniform illumination condition 

can be expressed by [9-6, 9-7] 

J (V ) ≅ J 0 (exp q(V − Rs J ) 

−1) +

V − Rs J − J ph (1) nkT Rsh 

where Jph, Jo, n, Rs and Rsh are the photocurrent density, saturation current density, diode ideality 

factor, series resistance, and shunt resistance, respectively. 

To study the effect of pulsed NLA treatment on the diode characteristics, photo- and 

dark- J-V measurements were performed at room temperature for each the selected cell. The 

measured parameters and calculated values of Jo, n, and Rs [9-8, 9-9] are summarized in Table 9­

4 (note that Rsh is quite high in all tested CIGS cells). 

From the dark- J-V curves shown in Figure 9-3, both the dark current density and 

saturation current density, Jo, of the NLA treated cells are consistently smaller as compared to 

the control cell. These curves suggest that the recombination current through surface defects 

dominates the dark- J-V curve under forward bias condition before NLA treatment (with diode 

ideality factor n >4). After the NLA treatment, a significant reduction in the dark current density 

was observed apparently due to reduction of surface defect density, and the recombination 

current was dominated by bulk defects in the junction space charge region at higher bias (n ≈ 2). 

These results indicate that defects in the surface region and interface of absorber and buffer 

layers can be effectively reduced by NLA treatment, thus decreasing the recombination of 

minority carriers via interface states and defects in the junction space-charge region of 

CdS/CIGS cells. Although cells treated by the NLA exhibit higher series resistance (Rs), the 

negative effect of Rs increase on the diode characteristic remains relatively small due to the 

extremely low value of Jo. 

Finally, the quantum efficiency (Q-E) was measured on the control and NLA treated 

cells. The results shown in Figure 4 indicate that for incident light with wavelengths greater than 

480 nm, the Q-E of the NLA cells is higher than that of the control cell, indicating that the NLA 

treatment improves the effective carrier lifetime and diffusion length in the absorber layer and 

hence increases the short-circuit current density in comparison with the control cell. In the 

shorter wavelength regime (λ< 0.48 μm), no improvement of Q-E was observed between the 

control cell and the NLA treated cells. It is also noted that values of the Q-E for cell annealed 

with 10 cycles of laser pulse were found to be lower than that of the cell with 5 cycles of laser 

pulse at same energy density, again suggesting that the optimal laser pulse number is between 5 

and 10 cycles. 

202
Table 9-4 Device performance parameters determined from photo- J-V 

measurements and device parameters calculated from dark- J-V measurements for 

the control cell and selected NLA-treated CIGS cells 

Cell # Ctrl-D-C3 CIGS-D1-C3 CIGS-D2-C3 

NLA condition None 30 mJ/cm2

, 5 pulses 30 mJ/cm2, 10 pulses 

Cell area (cm2

) 0.429 0.429 0.429 

Voc (V) 0.528 0.577 0.572 

Jsc (mA/cm2

) 29.78 34.24 32.00 

F.F. (%) 48.86 67.88 66.78 

Efficiency (%) 7.69 13.41 12.22 

Vm (V) 0.365 0.458 0.453 

Jm (mA/cm2

) 21.37 28.99 26.88 

N ~4.13 ~1.98 ~1.96 

J0 (mA/cm2

) ~3.22x10-3 ~1.1x10-3 ~1.06x10-3 

Rs (Ω) ~10.17 ~14.06 ~15.94 

Figure 9-3. Dark- J-V curve (semi-log plot) comparing the control cell to two 

NLA treated cells. 

203 
Figure 9-4. Quantum efficiency (Q-E) versus wavelength comparing the control cell to the two 

NLA-treated cells. 

The results of pulsed NLA treatment on the film properties and the performance of CIGS 

solar cells are very encouraging. Examination of the structural and electrical properties of films 

from 2 different sources clearly showed the annealing step modified the near surface region. 

Increased surface feature size evidenced in SEM photos and CIGS diffraction peak narrowing is 

consistent with increased crystallinity. Furthermore, DBOM measurements of the effective 

carrier lifetime indicated the lifetime could be increased by as much as a factor of 2.75. Hall 

measurements of CIGS samples deposited on SLG revealed NLA treatment increased the 

mobility and resistivity and decreased the net free hole concentration of each sample, consistent 

with the hypothesis of annealing out electrically active defects in the near surface region. Based 

on a parametric study, the best NLA result was obtained with a pulsed laser energy density of 30 

mJ/cm2

 and 5 pulse cycles. J-V and Q-E measurements were also made to study the effect of 

NLA treatment at the best condition on the performance of CIGS cells. The results show that 

pulsed NLA treatment has a beneficial effect on the cell performance with the cell efficiency 

increasing from 7.69 % to 13.41 and 12.22 % after annealing 2 different samples prior to device 

processing. The energy density of the laser beam and the number of pulse cycle were found to 

play a key role in changing the optical and electrical properties of the CIGS films and hence the 

cell performance. 

204
9.2 References cited 

[9-1] J. S. Luo, W. T. Lin, C. Y. Chang, W. C. Tsai, J. Appl. Phys., Vol. 82, p. 3621 (1997). 

[9-2] D. R. Chen, J. S. Luo, W. T. Lin, C. Y. Chang, P. S. Shih, Appl. Phys. Lett., Vol. 73, p. 

1355 (1998). 

[9-3] X. M. Lu, J. S. Zhu, W. S. Hu, Z. G. Liu, Y. N. Wang, Appl. Phys. Lett., Vol. 66, p. 2481 

(1995). 

[9-4] Sheng S. Li, B. J. Stanbery, C. H. Huang, C. H. Chang, Y. S. Chang, T. J. Anderson, 

“Effects of buffer layer processing on CIGS excess carrier lifetime: application of dualbeam optical modulation (DBOM) to process analysis [of solar cells],” 25th IEEE PVSC, 

Washington D.C., U.S.A., May 13-17, pp. 821-824 (1996). 

[9-5] C.H. Huang, Sheng S. Li, B.J. Stanbery, C.H. Chang, T.J. Anderson, “Investigation of 

buffer layer process on CIGS solar cells by dual beam optical modulation technique,” 

26th IEEE PVSC, Anaheim, U.S.A., Sep 29-Oct 3, pp. 407-410 (1997). 

[9-6] A. Fahrenbruch, R. Bube, Academic Press, New York, (1983). 

[9-7] S.M. Sze, Wiley, New York, (1985). 

[9-8] T. Yanagisawa, T. Kojima, T. Koyanagi, K. Takahisa, D. Nakamura, Elect. Lett., Vol. 36, 

pp. 1659-1660 (2000). 

[9-9] M.K. El-Adawi, I.A. Al-Nuaim, Vacuum, Vol. 64, pp. 33-36 (2002). 

205
10 Device Modeling and Simulation of CIGS Solar Cells 

10.1 Accomplishments 

Device modeling and simulation studies of a Cu(In1-x,Gax)Se2 (CIGS) thin film solar cell 

have been carried out. A variety of graded band-gap structures, including space charge region 

(SCR) grading, back surface region grading, and double grading of the CIGS absorber layer, are 

examined. The device physics and performance parameters for different band-gap profiles were 

analyzed. Based on the simulation results, an optimal graded band-gap structure for the CIGS 

solar cell is proposed. The performance of the optimally graded band-gap cell is superior to that 

of the uniform band-gap cell. The SCR grading of the CIGS absorber layer improves the opencircuit voltage (Voc) without significantly sacrificing the short-circuit current (Jsc) compared to 

the uniform band-gap CIGS. The back surface grading enhances both Voc and Jsc. An optimal 

graded band-gap profile, such as a double grading consisting of the SCR grading and back 

surface grading, improves significantly the efficiency up to 19.83% AM1.5G compared to the 

uniform band-gap profile with 15.42% efficiency. A comparison of the simulation results with 

published data for the CIGS cells shows an excellent agreement of photo- current densityvoltage and quantum efficiency characteristics. 

10.2 Introduction 

Much information is available in the literature regarding single-junction Cu(In1-x,Gax)Se2 

(CIGS) solar cells. Increasing the open-circuit voltage (Voc) to improve the overall performance 

of CIGS cells is highly desirable because it minimizes the interconnection losses in the 

manufacture of cell modules. When alloying the CuInSe2 (CIS) with Ga to form CIGS thin 

films, the wider band-gap energy of the CIGS absorber layer can potentially better match the 

solar spectrum, as well as increase the Voc of the fabricated cells, at the expense of a reduction in 

the value of the short-circuit current density (Jsc). Introducing a spatial variation of the Ga 

content within the CIGS layer, the band-gap profile can be optimized to increase the photon 

absorption and carrier diffusion. Furthermore, the Ga profile can be adjusted to optimize the 

CIGS absorber band gap profile, and hence improve the Voc and Jsc values. Thus, band-gap 

engineering geared to controlling the spatial distribution of the Ga content in the absorber layer 

can lead to enhancing the overall performance of CIGS cells. 

In this report, device modeling and numerical simulations of CIGS cells are conducted to 

analyze the impacts of various band-gap profiles for the CIGS layers on the performance 

parameters of the cells. Such modeling works will provide insight into which structures may be 

desired. The investigation focuses on the study of the effect on performance caused by various 

band-gap profiles in the space charge region (SCR) and in the back surface region of the 

absorber layer. Seven band gap profiles are analyzed, including the baseline case of a uniform 

band-gap profile. Based on the simulation results, an optimal graded band-gap structure for the 

CIGS cell is proposed. The simulation results of the optimized cell are compared to the reported 

experimental data. Additionally, the performance parameters of simulated results for CIGS cells 

with comparable grading structures in the absorber layers were found in excellent agreement 

with the published data. 

206
10.3Cell Structure and Material Parameters 

The schematic energy-band diagram under equilibrium condition for a typical 

ZnO/CdS/CIGS solar cell with a uniform band-gap profile is illustrated in Figure 10-1. An 

inverted surface layer, which is referred to as an ordered vacancy compound (OVC), is present 

between the CdS and CIGS layers in the diagram. Schmid et al. [10-1] have demonstrated the 

existence of such a thin OVC, In-rich n-type layer, on the surface of the as-deposited CIGS 

films, and assigned to it the stoichiometry of CuIn2Se3.5 or CuIn3Se5. However, some literatures 

[10-2] [10-3] revealed that the compound of In-rich surface layer might not be the OVC and used 

a more general term ‘surface defect layer’ for this thin n-type layer. The inverted surface layer is 

considered to be beneficial to the performance of CIGS cells because the electrical junction is 

shifted away from the high-recombination interface between the CdS and CIGS layer, and hence 

the recombination rate is reduced. However, the properties of the surface defect layer such as 

high defect densities and high resistivity might lead to deteriorating the cell performance. 

0.0 0.5 1.0 1.5 2.0 

Position (μm) 

Figure 10-1. The schematic energy-band diagram of a typical ZnO/CdS/CIGS solar cell under 

equilibrium condition. 

The CIGS cell structure considered in this study consists of the following material layers: 

n-ZnO, n-CdS, high-recombination interface, inverted surface, p-CIGS absorber, and a Mo on 

glass substrate. The computer simulation tool AMPS-1D (Analysis of Microelectronic and 

Photonic Structures) [10-4] was employed by specifying as input values the semiconductor 

parameters in each layer of the cell structure. The division of the layers for the cell structure is 

limited to the simplified device structure described above because of the limited knowledge 

available on the semiconductor parameters in each layer and uncertainties in the interface and 

junction properties arising from possible interdiffusion and reaction during the cell processing. 

The band gap of the CIGS absorber is engineered through the addition of a Ga profile. The 

anticipated changes in the physical properties of the CIS films with the addition of Ga include an 

increase in band gap, which mainly shifts the position of the conduction-band minimum [10-5], 

-9.0 

-8.0 

-7.0 

-6.0 

-5.0 

-4.0 

-3.0 

-2.0 

-1.0 

ZnO 

CdS 

Inverted surface layer 

Cu(In1-x,Gax)Se2 Eg 

Ec 

Ef 

Space charge region Ev

Energy (eV) 

207
as well as changes in the hole concentration [10-6], bulk defect densities [10-7], absorption 

coefficients, and electron affinities. The room temperature mobility was found to remain nearly 

constant while varying the Ga content over a wide range [10-6]. 

Except when otherwise mentioned, in the simulation studies an inverted surface layer 

with a thickness t = 30 nm, electron mobility μn = 10 cm2

/V-s, net carrier density n = 1012 cm-3, 

and band gap Eg = 1.3 eV is inserted between the CdS and CIGS layers. Also, an interface layer 

with a high density of effective recombination centers is placed at the metallurgical junction 

between the inverted surface layer and the CdS layer. The total thickness of the absorber layer is 

maintained at 2 μm for all cases, including the various graded band-gap structures studied. In 

order to model an effective recombination center, a deep level defect is placed in the middle of 

the band gap of (i) the CdS layer, (ii) the high-recombination interface layer, (iii) the inverted 

surface layer, and (iv) the SCR of the CIGS absorber. To examine the effects of band-gap 

gradient on the device performance, most of the material parameters used in the simulations are 

maintained unchanged, except that the band gaps, electron affinities, and optical absorption 

coefficients are varied with the Ga mole fraction in the CIGS films. 

10.4 Effect of Graded Band-gaps on Cell Performance 

First, a simulation study of uniform band-gap CIS and CIGS cells is carried out to 

establish a baseline for comparison. Two representative band-gaps, namely 1.04 eV for CIS and 

1.20 eV for CIGS, are considered. The simulation results are shown in Table 10-1. It is noted 

that the Cu(In1-x,Gax)Se2 cell with composition x ≈ 0.32 and Eg ≈ 1.20 eV [10-8] has a higher Voc 

and a smaller Jsc than the CuInSe2 cell corresponding to x ≈ 0 and Eg ≈ 1.04 eV, due to the 

smaller absorption coefficients and the reduction in the absorption of longer-wavelength photons. 

The effect of higher Ga-content of uniform band-gaps is not further analyzed because it has been 

reported in the literature that the device performance degraded when the Ga content in the 

Cu(In1-x,Gax)Se2 films exceeds 30 - 40%, i.e., when x > 0.3 - 0.4 [10-9]. 

Table 10-1. Simulated performance parameters of the Cu(In1-x,Gax)Se2 for x = 0 

(Eg = 1.04 eV), and Cu(In1-x,Gax)Se2 for x = 0.32 (Eg ≈ 1.20 eV) solar cells with a 

uniform absorber band-gap profile. 

x 

Eg 

(eV) 

η

(%) 

Voc 

(mV) 

Jsc 

(mA/cm2

) 

F.F. 

(%) 

CuInSe2 0 1.04 13.00 500 38.01 68.4 

Cu(In1-x,Gax)Se2 0.32 1.20 15.77 655 32.27 74.6 

Second, a simulation study of the dependence of the cell performance on the band gap in 

the SCR of CIGS absorbers is carried out. The recombination rate in the SCR, which controls 

the Voc of the CIGS cell, can be reduced by increasing the barrier height via the increase of the 

SCR band-gap. This improves the Voc, but a small loss of Jsc might be expected if the wider 

208 
band-gap region is not sufficiently thin. The dependence of Voc on the band gap in the SCR of 

CIGS cells has been reported [10-10]. In the simulation model, the SCR is composed of a series 

of layers of different band-gaps. The bottom-most layer in the SCR matches the band gap of 

1.04 eV of the underlying CIGS bulk material, and the subsequent SCR layers located above 

linearly increase until reaching a wide band-gap end-point value (EPV) varying from 1.16 to 

1.35 eV. More specifically, the EPV indicates the band gap of the SCR layer next to the inverted 

surface layer. The band-gaps of the inverted surface layer and of the remaining region of the 

CIGS bulk layer are maintained at Eg = 1.30 eV and 1.04 eV, respectively. The simulation 

results are shown in Figure 10-2, where the resulting performance parameters of Voc, Jsc, fill 

factor (F.F.), and efficiency (η) are shown as a function of the EPV of the band gap in the SCR. 

The value of Voc in Figure 10-2 (a) increases with increasing end-point band-gap in the SCR, 

improving from ~ 537 to ~ 551 mV, while the value of Jsc in Figure 10-2 (b) decreases slightly 

from ~ 37.9 to ~ 37.6 mA/cm2

. When comparing the performance of the graded band-gap cells 

reported in Figure 10-2 with the performance of the uniform band-gap CIS cell reported in Table 

10-1, it is apparent that SCR grading improves the Voc without significantly decreasing the Jsc. 

The F.F. values of the graded cells shown in Figure 10-2 (c) are also higher than that of the 

uniform band-gap CIS cell. The figure shows that different EPV values produce different F.F. 

values, but that the relationship is not monotonic. This is a consequence of the conduction-band 

offset between the inverted surface and CIGS layers. The value of η in Figure 10-2 (d) increases 

from ~ 15.3% to a maximum value of ~ 16.0% at an EPV value of ~ 1.30 eV. Comparing the 

efficiencies in Figure 10-2 (d) with the efficiency of the uniform-gap CIS cell in Table 10-1, it 

can be concluded that band-gap grading in SCR has a favorable effect on the final conversion 

efficiency. 

(c) (d) 

1.1 1.2 1.3 1.4 1.1 1.2 1.3 1.4 

End-point value (EPV) of band-gap in the space charge region (SCR) 

Figure 10-2. The dependence of cell performance parameters on the band-gap 

energy in the SCR of CIGS cells. 

Third, a simulation study was conducted to evaluate the effect of graded band-gap 

profiles of the CIGS absorber on the performance of the cells. A total of seven band gap profile 

cases are considered, as shown in Table 10-2. The simulation results for each case yield the 

performance parameters η, Voc, Jsc, and F.F reported in columns 3 to 6 of the table. The band525.0 

540.0 

555.0 

74.0 

77.0 

80.0 

14.5 

15.5 

16.5 

36.5 

37.5 

38.5 

(a) (b) 

V oc (mV) F.F. (%) 

J sc (mA/cm2 ) η (%) 

209
gap profile of Case 1 is the baseline, featuring a uniform band-gap of 1.16 eV for the entire CIGS 

absorber layer. In Case 2, a back surface field (BSF) is formed by the normal grading, where the 

band gap is linearly increased towards the back, and in Case 3, a step grading is used, where a 

wider band-gap layer is placed in the back. In Case 4, a wider band-gap CIGS layer of 1.30 eV, 

decreasing linearly to 1.16 eV, is placed entirely in the SCR. In Case 5 and 6, a double graded 

band-gap structure using the wider band-gap layer in the SCR and in the back surface region of 

the absorber is introduced. Comparing the performance of the cell in Case 1, both the Voc and Jsc 

with either a normal (Case 2) or step (Case 3) back surface grading are improved simultaneously, 

and hence the conversion efficiency is also increased, as shown in Table 10-2. The double 

graded band-gap structure of Case 6 realizes a high conversion efficiency of 18.39% AM1.5G. 

Figure 10-3 illustrates the schematic energy band diagram under equilibrium condition for Case 

6 with a double band-gap grading. In this figure, a quasi-electrical field [10-11] can be 

established by the band-gap gradient with the incorporation of a high Ga content in the back 

region (near the Mo contact) of the CIGS absorber layers by reducing the back surface 

recombination and increasing the effective minority-carrier diffusion length. This results in an 

efficient carrier collection in the CIGS cells, and thus both the Voc and Jsc are enhanced. The 

same trend was observed when comparing the performance of the cells in Cases 5 and 6 with that 

in Case 4. Case 7 is analogous to Case 6, except that the doping density in the CdS layer is 

increased to 6 × 1017 cm-3 and the deep- level defect densities are assumed to be equal to zero in 

the CIGS absorber. The result yields a maximum conversion efficiency of 19.83% for this cell. 

210
Table 10-2. Simulated performance parameters of the CIGS solar cells with 

various absorber band-gap profiles. 

η Voc Jsc F.F. Band-gap profile (%) (mV) (mA/cm2

) (%) 

2.0 μm 

Case 1 15.42 617 34.05 73.4 

E = 1.16 eV g 

1.1 μm 

Case 2 15.72 622 34.61 73.1 

1.16 eV 1.40 eV 

0.6 μm 

Case 3 15.81 620 34.77 73.3

1.16 eV 1.30 eV 

0.12 μm 

Case 4 17.57 658 33.55 79.6 

1.30 eV 1.16 eV 

0.12 μm 0.6 μm 

Case 5 18.34 669 34.46 79.6

1.16 eV 

1.30 eV 

1.30 eV 

0.12 μm 1.1 μm 

Case 6 18.39 675 34.31 79.4

1.30 eV 

1.40 eV 

1.16 eV 

Case 7 Same profile as Case 6 19.83 688 34.42 83.8 

211
-9.0 

-8.0 

-7.0 

-6.0 

-5.0 

-4.0 

-3.0 

-2.0 

-1.0 

ZnO 

CdS 

Inverted surface layer 

Cu(In1-x,Gax)Se2 

Ec 

Ef 

Ev

Energy (eV) 

Space charge 

region grading 

Back surface 

region grading 

0.0 0.5 1.0 1.5 2.0 

Position (μm) 

Figure 10-3. The schematic energy-band diagram of a CIGS cell with a double 

band-gap grading profile under equilibrium condition. 

In summary, a proper band-gap grading consisting of a wider band-gap CIGS layer in the 

SCR and a back surface grading are capable of significantly improving the performance of the 

CIGS cells. In the simulation studies the band gaps in the SCR and in the back region and the 

thickness of the back surface grading layers are varied to achieve the optimal performance 

through reducing the carrier recombination and increasing the carrier collection for the given 

material parameters and device structure. Comparing the performance of a uniform band-gap 

cell (Case 1) with the graded band-gap cells (Case 5 and Case 6), the benefit of band-gap grading 

for the CIGS solar cells is clearly demonstrated. 

10.5 Comparison with Experimental Data 

Attempts are made to compare the simulation results with reported experimental data for 

the CIGS cells. The measured values are taken from the photo- current density- voltage (J-V) 

and normalized quantum efficiency (Q.E.) curves of an experimental CIGS cell with 18.8% 

efficiency under standard conditions produced at the National Renewable Energy Laboratory 

(NREL) [10-12] [10-13]. The performance of this experimental cell is compared with the 

simulated results of Cases 6 and 7 in Table 10-2, and the results are shown in Figures 10-4 and 

10-5, where the solid line represents the simulated cell of Case 6, the dashed line represents the 

simulated cell of Case 7, and the solid circles denote the experimental values. Figure 10-4 shows 

that the photo- J-V curve for Case 6 closely matches the experimental curve for the NREL CIGS 

cell. The photo- J-V values obtained from the simulation of Case 7 suggest that a higher 

efficiency is possible. Figure 10-5 shows a comparison of the simulated Q.E. curve for Case 6 

with the experimental curve. Good agreement was also observed in this case in the wavelength 

(λ) range of 0.5 μm ≤ λ ≤ 1.2 μm. 

212
J (mA/cm2 ) 

0.0 

5.0 

10.0 

15.0 

20.0 

25.0 

30.0 

35.0 

40.0 

0.0 0.2 0.4 0.6 0.8 

V (V) 

Figure 10-4. The photo-current density-voltage (J-V) curves for the simulated 

results (solid and dashed lines) and the reported experimental data (solid circles) 

[10-12] [10-13]. Normalized Q.E. 

0.0 

0.2 

0.4 

0.6 

0.8 

1.0 

0.5 0.6 0.7 0.8 0.9 1.0 1.1 1.2 1.3 

λ (μm) 

Figure 10-5. The normalized QE for the simulated results (solid line) and the 

reported experimental values (solid circles) [10-12]. 

A separate analysis is carried out to fit the performance of the simulated parameters with 

data available in the published literature for CIGS cells. Dullweber et al. have reported 

experimental results for band-gap grading in CIGS cells [10-10] [10-14]. The band-gap structure 

of the CIGS absorber layer in Table 10-3 is based on the Auger depth profile reported by 

Dullweber et al. [10-10], which allows defining the layer dimensions and compositions for the 

simulation model. For the simulation study, the depth profile is represented by an equivalent 

213
band-gap profile in the CIGS absorber using the equation Eg (y) [eV] = 1.02 + 0.67 y +

0.11 y (y-1) proposed in [10-15], where the variable y denotes the ratio of Ga/(Ga+In). The 

resulting band-gap profile reveals a double-graded CIGS absorber layer with linearly graded 

band-gap structures in the SCR and in the CIGS bulk layers, as shown in Table 10-3, and is used 

as a set of input parameters for the simulation study. This notch structure is a particular case of 

double grading. Theoretical analysis and simulation studies of this structure have been reported 

[10-16]. The measured data are also taken from the J-V curve given in [10-10]. The key 

parameters used in fitting the simulation results to the experimental values are the defect density 

in the inverted surface layer, the defect densities of the mid-gap deep- level in the SCR, and the 

defect and hole densities in the CIGS absorber. The results are shown in Figure 10-6, where the 

simulated values are plotted using a solid line, and the experimental values are indicated by the 

solid circles. Although the simulated F.F. is slightly higher than the experiment, excellent 

agreement was observed in the J-V curve. 

Table 10-3. The energy band-gap profile of the CIGS absorber layer used in 

simulation. 

1.70 μm 

0.16 μm 

1.23 eV 1.07 eV 1.45 eV 

0.0 0.2 0.4 0.6 0.8 

V (V) 

Figure 10-6. The photo-current density-voltage (J-V) curves for the simulated results (solid line) 

and the published data (solid circles) [10-10]. J (mA/cm2 ) 

0.0 

5.0 

10.0 

15.0 

20.0 

25.0 

30.0 

35.0 

40.0 

214
Finally, another work by Dullweber et al. [10-14] for CIGS cells with a normal or reverse 

graded absorber is considered for fitting the simulation results to experimental values. An 

absorber with normal grading profile is one where the band gap is increased linearly from the 

front to the backside in the entire absorber, while the absorber with reverse grading profile is one 

where the band gap is decreased linearly from the front to the back. To investigate these linearly 

graded structures via simulation, the composition depth profiles of the CIGS absorber given in 

[10-14] were used for the calculation of the band-gap profiles for the CIGS absorber. From the 

represented band-gap profiles, an absorber structure with normal grading is increased linearly 

from the front band-gap of 1.03 eV to the backside band-gap of 1.22 eV, while an absorber 

profile with reverse grading is decreased from 1.24 eV to 1.06 eV. Also, a uniform band-gap 

profile of 1.14 eV is considered. Experimental performance parameters are also extracted from 

[10-14] for the cases of normal, reverse, and uniform grading, and are compared with the 

simulated results. The major parameters used in fitting the simulated results to the experimental 

values are as described in the preceding paragraph. The results are shown in Figure 10-7, where 

the solid circle and triangle denote the simulated values and experimental data, respectively. The 

x-axis indicates CIGS absorber band-gap profiles consisting of the reverse grading, uniform, and 

normal grading from left to right. It is noted that from Figure 10-7 (a) Voc is found to decrease 

from 565 to 509 mV (a change of 56 mV), from the reverse to the normal grading, while the 

value of Jsc increases from 32.5 to 37.5 mA/cm2

 (a change of 5 mA/cm2) as shown in Figure 10-7 

(b). These trends are also found in the experimental values. Excellent agreement was obtained 

on the performance parameters Voc, Jsc, and F.F. Even though the values of η in the simulation 

are slightly higher than their experimental counterparts, a reasonable fit is nevertheless obtained. 

215
0.4 

0.6 

0.8 

20 

40 

60 

50 

70 

90 

8 

13 

18 

V oc (V) J sc (mA/cm2 ) F.F. (%) η (%) 

(a) 

(b) 

(c) 

(d) 

Reverse grading Uniform Normal grading 

Figure 10-7. A comparison of the performance parameters for the simulated 

results (circles) and the published data (triangles) [10-14]. 

10.6 Conclusions 

Comprehensive device modeling and numerical simulation studies for single-junction 

CIGS thin film solar cells have been carried out using the AMPS-1D device simulation program, 

including various absorber band-gap profiles. A detailed analysis of the graded band-gap 

structures for the CIGS cells has been presented. It is shown that band-gap grading using a wide 

band-gap layer in the SCR and/or back surface grading using the gradient of Ga content in the 

CIGS absorber layer greatly improve the performance of single-junction CIGS cells. 

Additionally, studies to fit the simulation results with the reported experimental data show 

reasonable agreement in the J-V and Q.E. curves between the modeling and the experimental 

data. 

10.7 References Cited 

[10-1] D. Schmid, M. Ruckh, F. Grunwald, and H.W. Schock, “Chalcopyrite/Defect 

Chalcopyrite Heterojunctions on the Basis of CuInSe2”, Journal of Applied Physics, Vol. 

73, No. 6, pp. 2902-2909 (1993). 

216
[10-2] R. Herberhoz, U. Rau, H.W. Schock, T. Haalboom, T. Gödecke, F. Ernst, C. Beilharz, 

K.W. Benz, and Cahen, ”Phase Segregation, Cu Migration and Junction Formation in 

Cu(In,Ga)Se2“, European Physical J. Appl. Phys, Vol. 6, pp. 131-139 (1999). 

[10-3] J. Guillemoles, T. Haalboom, T. Gödecke, F. Ernst, D. Cahen, ”Phase and Interface 

Stability Issues in Chalcopyrite-based Thin Film Solar Cells”, Mat. Res. Soc. Symp. 

Proc., 485, pp. 127-132 (1998). 

[10-4] H. Zhu, A.K. Kalkan, J. Hou, and S.J. Fonash, “Application of AMPS-1D for Solar Cell 

Simulation”, AIP Conference Proceedings, No. 462, pp. 309-314 (1999). 

[10-5] S.H. Wei and A. Zunger, ”Band Offsets and Optical Bowings of Chalcopyrites and Znbased II-VI Alloys”, J. Appl. Phys., Vol. 78, pp. 3846-3856 (1995). 

[10-6] D.J. Schroeder, J.L. Herberholz, and A.A. Rockett, “Point Defects and Hole Transport in 

Epitaxial CuIn1-xGaxSe2”, 11th Int. Conf. on Ternary and Multinary Compounds, pp. 749­

752 (1999). 

[10-7] G. Henna, A. Jasenek, U. Rau, and H.W. Schock, ”Influence of the Ga-content on the 

Bulk Defect Densities of Cu(In,Ga)Se2”, Thin Solid Films, Vol. 387, pp. 71-73 (2001). 

[10-8] S.H. Wei, and S.B. Zhang, and A. Zunger, “Effects of Ga Addition to CuInSe2 on Its 

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218
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Subcontract Report 

3. DATES COVERED (From - To) 

2 October 2001–30 September 2005 

4. TITLE AND SUBTITLE 

Fundamental Materials Research and Advanced Process 

Development for Thin-Film CIS-Based Photovoltaics: Final 

Technical Report, 2 October 2001 – 30 September 2005 

5a. CONTRACT NUMBER 

DE-AC36-99-GO10337 

5b. GRANT NUMBER 

5c. PROGRAM ELEMENT NUMBER 

6. AUTHOR(S) 

T.J. Anderson, S.S. Li, O.D. Crisalle, and V. Craciun 

5d. PROJECT NUMBER 

NREL/SR-520-40568 

5e. TASK NUMBER 

PVB65101 

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University of Florida 

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National Renewable Energy Laboratory 

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14. ABSTRACT (Maximum 200 Words) 

The objectives for this thin-film copper-indium-diselenide (CIS) solar cell project cover the following areas: Develop and 

characterize buffer layers for CIS-based solar cell; grow and characterize chemical-bath deposition of Znx Cd1-xS buffer 

layers grown on CIGS absorbers; study effects of buffer-layer processing on CIGS thin films characterized by the dualbeam optical modulation technique; grow epitaxial CuInSe2 at high temperature; study the defect structure of CGS by 

photoluminescence spectroscopy; investigate deep-level defects in Cu(In,Ga)Se2 solar cells by deep-level transient 

spectroscopy; conduct thermodynamic modeling of the isothermal 500°C section of the Cu-In-Se system using a defect 

model; form α-CuInSe2 by rapid thermal processing of a stacked binary compound bilayer; investigate pulsed non-melt 

laser annealing on the film properties and performance of Cu(In,Ga)Se2 solar cells; and conduct device modeling and 

simulation of CIGS solar cells. 

15. SUBJECT TERMS 

PV; materials research; thin film; advanced process development; solar cell; copper indium diselenide; chemical-bathdeposition; epitaxial growth; optical modulation; photoluminescence spectroscopy; deep-level transient spectroscopy; 

thermodynamic modeling; simulation; CIGS. 

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