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| **PROCESSING MATERIALS, DEVICES AND DIAGNOSTICS FOR THIN FILM         |
| PHOTOVOLTAICS:**                                                      |
|                                                                       |
| **[FUNDAMENTAL AND MANUFACTURABILITY ISSUES]{.smallcaps}**            |
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| **Annual Report**                                                     |
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| **to**                                                                |
|                                                                       |
| **National Renewable Energy Laboratory**                              |
|                                                                       |
| **under**                                                             |
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| **Subcontract No. ADJ-1-30630-12**                                    |
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| **9/05/02 to 9/04/03**                                                |
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![](media/image1.wmf){width="2.611111111111111in"
height="1.1388888888888888in"}

**Institute of Energy Conversion**

**United States Department of Energy**

**University Center of Excellence**

###### For Photovoltaic Research and Education

(National Renewable Energy Laboratory)

Annual Report

to

National Renewable Energy Laboratory

under

Subcontract No. ADJ-1-3-630-12

9/05/02-9/04/03

R. W. Birkmire, W. N. Shafarman, E. Eser,

S. S. Hegedus, B. E. McCandless, R. Aparicio, K. D. Dobson

[Institute of Energy Conversion]{.smallcaps}

[University of Delaware]{.smallcaps}

[United States Department of Energy]{.smallcaps}

[University Center of Excellence]{.smallcaps}

[For Photovoltaic Research and Education]{.smallcaps}

[(National Renewable Energy Laboratory)]{.smallcaps}

[Newark, Delaware 19716-3820]{.smallcaps}

# 

# **SUMMARY**

##### CdTe/CdS based solar cells

Investigations of the dynamics of CdTe etching have continued. Variable
Angle Spectroscopic Ellipsometry (VASE) has shown that bromine-methanol
(BM) etching to be consistent and very reproducible. BM polishing of
CdTe will be used for sample preparation for a systematic investigation
of nitric-phosphoric (NP) etching of CdTe, to accurately determine Te
thickness.

A number of devices, contacted with Cu-containing graphite paste, were
prepared to monitor the effects of different etch treatments on cell
performance and stability. Devices processed from Br~2~-based etches of
high viscosity, and hence extremely thin Te-rich surface films, were
found to shunt within a few hours of processing, due to the inability of
the thin Te to getter Cu. Devices processed using NP etches of different
times, and contacted with Cu-containing graphite paste, initially all
showed similar performance. Following stressing, no decrease of V~OC~ or
J~SC~ was observed, however, the R~OC~ of devices that had received
shorter NP etches had significantly increased, while longer etched
samples showed very little change in R~OC~. These changes are indicative
of back contact degradation and again suggest that devices with back
contacts processed with thicker Te components will exhibit improved
stability due to improved Cu gettering and compensation of other
chemical changes.

The chemistry of Cu with Te-rich CdTe surfaces, prepared by different
etch treatments, was investigated. The Cu was supplied from graphite
paste contacts. Differences in the copper telluride product were
observed, dependent on etch type and chemical nature of Cu. The thicker
Te-rich layer formed by NP etching results in a higher degree of the
CuTe phase, on reaction with Cu-powder or Cu salts in the graphite
paste, compared to BM etching. CuCl and CuI appear to form Cu~2-x~Te on
BM etched CdTe, while Cu powder forms CuTe. Initial results were
complicated by the presence of residual crystalline graphite from the
paste, however this residue can be easily removed.

Effects of stressing cell components and completed devices in conditions
of varying humidity were monitored. TeO~2~ and CdTeO~3~ were found to be
formed at lower humidities on CdTe, while the formation of amorphous
hydroxide, oxide or hydrate phases are more likely at higher humidities.
The IEC back contact appears to be stable at higher humidities, but is
oxidized to TeO~2~ at low humidity. This suggests that copper tellurides
are more susceptible to oxidation by atmospheric oxygen. However,
despite the extensive chemical changes observed, device performance was
almost unaffected. In contrast, First Solar devices showed significant
degradation. The robustness of the IEC back contacts is proposed to be
due to the thick, 100 nm, Te film present in the back contact, which
compensates chemical changes to the contact.

High efficiency CdTe/CdS devices require a robust and durable thin CdS
film, which maintains a continuous layer for uniform junction formation.
We investigated the relationship between device processing, materials
properties and device performance for cells with CdS films deposited by
chemical surface deposition (CSD), which produces conformal ultra-thin
CdS films with low particulate density and high utilization of Cd
species. In CSD, a solution at ambient temperature is applied to a
pre-heated glass/TCO substrate. Film deposition proceeds by
heterogeneous nucleation at the surface as heat is transferred to the
solution. Heat loss from the solution to the ambient helps maintain
conditions favorable for film growth over the time needed typically \~ 5
min. The focus has been on establishing and understanding baseline
processing procedures for the CdS/TCO superstrates and post-deposition
procedures for completed CdTe/CdS structures. The processing variables
investigated here for CdS/TCO were time and temperature of the CdCl~2~
treatment given to the CSD CdS prior to CdTe deposition and use of a
high temperature anneal (HTA) step. Measurements of a film 90 nm thick
CdS films on Ga~2~O~3~/SnO~2~/glass before and after vapor CdCl~2~
treatment at 430°C shows significant recrystallization based on
sharpening of the transmission band-edge and sharpening of diffraction
peaks (GIXRD). It was found that CdS films grown on Ga~2~O~3~ have
larger grain size than those on glass.

Having established that CdCl~2~ treatments change the CdS properties, we
optimized the device performance by varying CdCl~2~ temperatures
(350-430°C) and times (5-20 minutes). In general, efficiency improved
with increasing the temperature and time. Therefore, a temperature of
430°C and time of 20 min was selected for CdS treatment in subsequent
studies.

Since the CdS is deposited from a Cl based-solution, it was possible
that a HTA would be sufficient to activate the CdS grain growth against
further interdiffusion, thus, eliminating the need for a separate
CdCl~2~ treatment of the CdS. Devices were made with and without the HTA
and CdCl~2~ treatments to CdS prior to CdTe deposition. The CdS CdCl~2~
treatment was at 430°C for 20 min and the HTA was at 550° for 30 min in
Ar. Also included in this study were substrates of glass/SnO~2~/CdS from
First Solar. They did not receive any treatments since the CdS was
deposited at high temperature. J-V results showed that cells with the
CdS HTA alone or in combination with CdCl~2~ had reduced FF. The best
cell on three identically processed pieces with CdCl~2~ only (no CdS
HTA) ranged from 10-11%. Yields were better than 75%. Cells with HTA had
9-10% efficiency. Devices on First Solar CdS had similar V~OC~ and FF to
cells on CSD CdS. We conclude that our CSD CdS is not a major limitation
to higher V~OC~ or yield at this time providing it has a CdCl~2~
pre-treatment.

We are continuing to investigate the relationship between CdTe
deposition and film properties for cells with CdTe deposited at a high
rate on moving substrates by vapor transport (VT). The VT deposition
process is similar but not identical to that used at First Solar. The
influence of He and O~2~ flow rate (2% O~2~/He) and pressure during the
CdTe growth was investigated using AFM, GIXRD and contact wetting angle.
As compared to the baseline run, increasing the total flow of helium
through the source, results in a substantial decrease in grain size.
Decreasing the total flow of He yields slightly larger, more densely
packed CdTe grains. Increasing the ratio of helium to oxygen by two,
results in larger less densely packed grains. Device results from these
runs are pending.

The relation between device processing and device performance for cells
with VT CdTe deposited at a high rate on moving substrates has focused
on establishing and understanding baseline processing for the
post-deposition procedures of completed CdTe/CdS structures. For the
CdTe/CdS post-deposition treatments, the time and temperature of the
CdCl~2~ treatment, use of a HTA prior to CdCl~2~ treatment, and spacing
between the CdCl~2~ plate and CdTe surface were investigated. Following
deposition of a \~4 µm CdTe layer at 550°C in a single 6 minute pass in
He/O~2~ at 20 T, pieces receive a HTA and/or CdCl~2~ treatment. Contact
processing consists of a BDH etch then loading the samples into the
evaporator for the layered back contact structure of 50 nm Te, 30 nm Cu,
50 nm Au. Post-contacting annealing was investigated, and a 30 minute
anneal in Ar at 190°C was found to be optimum.

The time and ambient of the CdCl~2~ treatment at 430°C of the completed
CdS/CdTe structure was varied from 0 to 60 minutes in either dry or
moist air. It was found that the HTA by itself (no CdCl~2~ treatment)
accomplished the same enhancement in performance, to \~8% efficiency, as
the optimum CdCl~2~ treatment in this experiment. A device with neither
HTA nor CdCl~2~ had poor FF and poor collection from the CdTe. There was
no clear difference in device performance between dry or moist air
during the CdCl~2~ treatment, suggesting the BDH treatment rendered the
CdTe surfaces equally clean in both cases. However, before BDH there
were differences in the type of oxides detected by GIXRD and in the
surface energy determined by the contact-wetting angle.

It has been shown that O~2~ is critical for the CdCl~2~ treatment to
achieve full impact on the device. This leads to concerns about
sufficient circulation and replenishment of the O~2~ ambient during the
CdCl~2~ treatment in a static system. Humidity and exposure to high
temperatures in presence of air can lead to oxide formation on the
CdCl~2~ source plate, leading to concern about their reuse. To study
these variables, the spacing between the CdCl~2~ plate and the substrate
was varied during the CdCl~2~ treatment. Increasing the distance would
increase the availability of O~2~ at the CdTe surface. Cells made with
fresh plates included those with 0 and 6X spacing where X is our
standard spacing of \~0.6 mm. After being used for treating the cells
with no spacing, the CdCl~2~ plate was re-used for treating with 2X
spacing. It was found that the largest spacing (6X) was best,
demonstrating that the delivery of O~2~ and CdCl~2~ are essential to
obtaining good device performance. Devices processed with used CdCl~2~
plates had much lower efficiency and/or severe retrace instability.
Following this result, fresh CdCl~2~ plates were used for every
treatment. The best device from this series was VT85.33-01, having
V~OC~=0.749 V, J~SC~=22.9 mA/cm^2^, FF=65.2%, and eff.=11.2%.

**CuInSe~2~-based Solar Cells**

Development of the roll-to-roll deposition of CIGS on polyimide web
focused on film cracking and / low adhesion with specific effort on
finding a correlation with the Mo deposition process. Mo films were
deposited onto two types of Upilex (S and SG) with and without oxygen
plasma treatment and at different temperatures and gas pressures but
CIGS film cracking and adhesion were found to be insensitive to these
variations. A number of changes were made to the CIGS roll-to-roll
deposition system to improve its robustness and reproductibility. In
this respect, a new Se source design now allows the control of the
surface temperature of the liquid Se source. Shielding of the
thermocouple extension wires made possible the temperature control of
the source effusion rates, which is more robust than the Atomic
Absorption Spectroscopy control.

Work on CIS-based devices with increased V~oc~ focused on adhesion and
material and device performance of Al containing alloys, and on S
diffusion into CIS and CIGS films. Al alloy films have been deposited by
1-step (constant flux), 2-step (Al-rich flax followed by Cu-free flux),
and 3-step (Cu-free flux at low temperatures followed by only Cu flux
and then Cu-free flux at higher temperatures processes. Going from
1-step to 3-step processes grain size was found to increase. Film
adhesion was similar for 1-step and 2-step processes, but much worse for
the films prepared by the 3-step process. Device results showed similar
efficiencies for 1-step and 2-step processes, decreasing with increasing
Al content. Films prepared by the 3-step process could not be made into
devices except for the lowest Al content films due to the poor adhesion.
QE data showed increasing voltage dependent collection with increasing
Al content, i.e. increasing bandgap.

Characterization of the physical vapor deposited films on Cu, Al and Se
to form the ternary CuAlSe~2~ showed different structure depending on
the ratio of Se flux R~SE~ to total metal flux R~M~. R~SE~/R~M~ 910
gives mixed phases of Cu-Al inter-metallic phases plus Al~2~Se~3~ to
CuAlSe~2~ + Cu~2~Se going through single-phase CuAlSe~2~ depending on
the Cu/Al ratio. For R~SE~/R~M~ \>10 the structure is a mixture of
CuAlSe~2~ and a new phase β-CuAlSe~2~.

In the area of fundamental studies, alternative emitter layer materials
and processes were evaluated and optical properties and electronic
transitions in Cu(InGa)Se~2~ films were determined through spectroscopic
ellipsometry. Air annealed chemical bath deposited ZnS films to form
mixed sulfide-oxide layers were characterized both optically and
structurally. While bandgap decreases with air anneal intensity,
glancing angle incidence XRD does not indicate any sulfide-oxide alloy
formation. Devices made with the ZnS window layer do not show any
correlation between efficiency and air-anneal time. CdS layers formed by
chemical surface deposition method were evaluated for their potential in
Cu(InGa)Se~2~ based devices. The process is highly desirable from a
manufacturing point of view, since it has high utilization rate and low
waste; devices made with CdS deposited by this process demonstrated high
efficiencies showing the expected promise.

Optical constants of polycrystalline thin film CuIn~1-x~Ga~x~Se~2~
alloys with Ga/(Ga+In) ratios from 0 to 1 have been determined by
spectroscopic ellipsometry over an energy range from 0.75 to 4.6eV.
CuIn~1-x~Ga~x~Se~2~ films were deposited simultaneous thermal
evaporation of elemental copper, indium, gallium and selenium. X-ray
diffraction measurements show that the CuIn~1-x~Ga~x~Se~2~ films are
single phase. Due to their high surface roughness, the films are
generally not suitable for ellipsometer measurements. An innovative
method was developed in which spectroscopic ellipsometer measurements
were carried out on the reverse side of the CuIn~1-x~Ga~x~Se~2~ Films
immediately after peeling them from the Mo-coated soda lime glass
substrates. A detailed description of multilayer optical modeling of
ellipsometric data, generic to ternary chalcopyrite films, is presented.
Accurate values of the refractive index and extinction coefficient were
obtained and the effects of varying Ga concentration on the electronic
transition were determined

##### Si-BASED SOLAR CELLS

The grain size enhancement of HWCVD Si films deposited on Al layers has
been investigated in two temperature regimes. These regimes are
determined by the eutectic temperature of 575°C in the Si-Al system.
With the formation of a liquid phase above the eutectic temperature, a
different mechanism is expected to govern the grain enhancement process
compared to that expected below this temperature where the Si-Al mixture
remains in the solid state. This report describes the results obtained
for depositions of HWCVD Si films on Al layers at 430°C and 600°C.

Aluminum layers with thickness varying from 0.01 to 1.0 μm were
deposited onto Corning 7059 type glass substrates using electron-beam
evaporation. Silicon layers were subsequently deposited onto the Al
coated samples in the HWCVD system. Silane was used as the source gas
with no hydrogen dilution. Silicon layer depositions were carried out at
a substrate temperature of 430°C and 600°C, filament temperature of
1850°C and pressure of 25 mTorr. The films were deposited at two
different growth rates of 1.0 and 5.0 μm/hr, which were attained by
varying the silane flow rate from 4.5 to 22.5 sccm, respectively. Films
with thickness values of 2 and 5 μm were deposited with deposition times
corresponding to 24 min and 5 hrs, respectively.

For deposition at 430°C, Raman spectroscopy and XRD showed Si films
deposited on Al had random orientation and higher degree of
crystallization compared to films deposited on glass. The crystalline Si
Raman profile of films deposited on glass had a typical tail extending
toward lower wavenumbers indicative of a-Si and nanocrystalline Si
phases. In contrast, films deposited on Al, had profiles similar to
those obtained for c-Si. For 2 μm thick films grown at 1 μm/hr, the
average grain size and crystalline fraction increases with decreasing
Si/Al ratio. The maximum average grain size of 200 nm was achieved with
a Si/Al ratio of 2. As the growth rate is increased from 1 to 5 μm/hr,
the average grain size and crystalline fraction decreases at a constant
Si/Al ratio. AFM analysis revealed that some areas on the film surface
are covered with small particles, 10-50 nm in diameter. It is speculated
that these features are individual grains, which do not extend
throughout the bulk of the Si film. TEM analysis found heterogeneous
structure of several large grains but mostly nano-crystalline regions.
These two crystalline regions can be found either side by side on the Al
layer or as layers with one evolving from the other. It is speculated
that the presence of the nano-crystalline phase indicates an incomplete
grain enhancement process and that the effect of lower growth rate for
given film thickness is to increase the interaction time between the
layers. The poly-Si does not form a continuous layer. The bulk of the
films also contain Al. Al layer inversion is incomplete, perhaps due to
insufficient interaction time between the Si and the Al layer. For 5.0
μm thick samples the average grain size varied between 40 and 70 nm, and
the crystalline fraction did not change significantly with the growth
rate. The decrease in grain size for Si films with 5 μm thickness may be
due to the existence of a limiting thickness beyond which the Al ceases
to have an effect on the growth process. If this hypothesis is correct,
the film evolves from a large grain to a nanocrystalline grain
microstructure as seen on films deposited on glass.

Si films deposited at 600°C exhibit larger average grain size as
measured by XRD peak broadening. This contrast is more pronounced as the
growth rate and Si film thickness increases. However, all films were
non-homogeneous in terms of morphology and composition perhaps as a
result of poor wetting of the Al layer on the 7059 glass at 600°C or
alternatively due to segregation of the liquid phase forming at the
eutectic composition. On average, the crystalline volume fraction of
films deposited at 600°C is similar to those deposited at 430°C.

Fabrication of p-i-n solar cells with all PECVD µc-Si layers was
investigated on glass/SnO~2~ substrates. The 2.0 µm i-layer was
deposited for 9 hours. No ZnO or a-Si buffer layers or any other
transitional layers were incorporated. Device performance was very poor,
with V~OC~=0.2V and efficiency \<1%. CV and QE measurements suggested a
relatively high defect density and low barrier height. The µc-Si devices
have substantially more absorption below the a-Si absorption edge
confirming that their structure is between a-Si and c-Si.

Thin film Si n-i-p solar cells were fabricated on n-doped c-Si
substrates using PECVD and HWCVD microcrystalline i-layers. No metal
induced crystallization was utilized. The structure of the devices was
Ag-Pd-Ti/n^+^-cSi/(PECVD or HWCVD) i-μcSi/PECVD p-μcSi/ITO/Ni-Al. The n+
wafer substrates served as the n-layer in the p-i-n cells. In addition,
diagnostic n-p devices were fabricated by depositing PECVD p-μcSi
emitter layers directly onto the n^+^-c-Si substrates. Identical n-i-p
and n-p devices were also fabricated which incorporated a 30 nm a-Si
buffer layer between the p-layer and i-layer or n+ c-Si. The performance
of all solar cell devices was quite poor. All cells with a HW or GD
i-layer absorber had less than 1% efficiency although many had well
formed diode characteristics in light and dark.

Both n-i-p and n+/p+ cells using GD and HW µc-Si absorber i-layers have
produced cells with poor photo-carrier collection and high
recombination. The a-Si buffer layer increases V~OC~ but reduces J~SC~
and FF in both n-i -p and n+/p+ cells. No first-order differences were
seen between devices the HW or GD µc-Si layers. n+/p+ cells with the
a-Si buffer achieved V~OC~=0.62 V but with very poor J~SC~ and FF.

Previously, standard a-Si p-i-n devices routinely achieved FF=70% while
recently FF has been limited to 50-60%. QE measurements indicated
poorhole collection from the i-layer. The most likely cause is
donor-like impurities such as P or O, which reduce the lifetime of the
hole. Considerable effort was made in cleaning or modifying the PECVD
deposition system to eliminate possible sources of O or P. None of these
changes resulted in a higher FF. Select samples were sent to NREL for
SIMS analysis. Since our PECVD reactor is a single chamber system, a
major concern in both a-Si and µc-Si devices is the P and B
cross-contamination. Results showed that P and B are quite low in PECVD
and HWCVD a-Si films and devices. Some recent devices have had B
concentrations of 10^18^ cm^-3^ in the p-layer while P was 10 times
higher. This suggests our burying layer is not effective anymore in
preventing P cross-contamination. Consequently, a new burying procedure
is being implemented. For PECVD or HWCVD µc-Si devices, the
concentration of O, N and C are several orders of magnitude higher than
those in the a-Si cell and the c-Si wafer, as is reported by others.
These impurities in the HWCVD samples have a concentration gradient,
which decreases from the top surface and level off to a constant bulk
concentration at approximately one micron into the film. It is
speculated that the lower H content in HWCVD films combined with the
excessive number of grain boundaries and film porosity may facilitate
contaminant penetration into the films.

Boron doped HWCVD films were deposited and characterized. Films were
polycrystalline with typical (220) preferred orientation. Crystalline
fraction was 75 to 90%. A dependence of conductivity with dopant gas
concentration and film thickness was found. For submicron films,
conductivies were constant at 10^-6^ S/cm for B~2~H~6~/SiH~4~ \<10^-3^.
At B~2~H~6~/SiH~4~ =10^-2^, the conductivity reached 10^-3^ S/cm. For
films with thickness \> 2 μm and B~2~H~6~/SiH~4~ =10^-2^, the
conductivity further increased to 10^-1^ S/cm. The effect of thickness
on conductivity points to a change in microstructure as the film
thickness increases. This hypothesis is corroborated by activation
energy data, which indicates a change in mobility with thickness. Future
work will include a determination of the effects of film microstructure
and dopant gas concentration on film conductivity.

##### DIAGNOSTICS

Contact wetting angle, ellipsometry and glancing incidence x-ray
diffraction are surface sensitive technique methods that are under
investigation as diagnostic tools for thin-film processing. An apparatus
for carrying out contact angle measurements was constructed and wetting
angle measurements are being carried out on different types of samples
to determine the sensitivity to changes in surface energy arising from
thin-film processing. Ellipsometry is being developed to determine film
composition in multi-layer thin-film structures such as treated CdTe
surfaces. Glancing x-ray diffraction is being developed to determine
residual stress in thin films and the thickness of surface layers
obtained on CdTe films under different post-deposition processing
conditions and to evaluate residual stress in as-deposited Mo films.

# **LIST OF CONTRIBUTORS**

## 

#### PRINCIPAL INVESTIGATORS

Roger Aparicio

Robert W. Birkmire

Kevin D. Dobson

Erten Eser

Steven S. Hegedus

Brian E. McCandless

William N. Shafarman

##### RESEARCHERS

Wayne A. Buchanan

##### VISITING PROFESSIONALS

Vivian Alberts, Rand Afrikaans Univ., S. Africa

**POST-DOCTORAL FELLOWS/**

**LIMITED TERM RESEARCHER**

Estela Calixto-Rodriguez

Mario Gossla

Greg Hanket

Sylvain Marsillac

Takashi Minemoto

Puthar Devassey Paulson

Sergey Rykov

Baosheng Sang

Udai Singh

##### TECHNICAL SUPPORT

Joshua Cadoret

Vicky DiNetta

Ronald Dozier

Shannon Fields

Kevin Hart

Thomas Hughes-Lampros

**GRADUATE STUDENTS**

Darshini Desai

Ozgenc Ebil

Moses Haimbodi

Toby Junker

Jochen Titus

Mike Zellner

##### STUDENTS

Michael Angelo

John Dellose

Jeffrey Doody

Jonathan Frey

Scott Kominsky

Mary Kondos

Kunal Modi

Laura Robinson

Daniel Ryan

##### DOCUMENT PREPARATION

Steven S. Hegedus

L. Bonnie Johnston

# TABLE OF CONTENTS

*SUMMARY [III](#summary)*

LIST OF CONTRIBUTORS [X](#list-of-contributors)

TABLE OF CONTENTS [XII](#table-of-contents)

LIST OF FIGURES [XIV](#list-of-figures)

1 Objectives [3](#objectives)

1.1 CdTe/CdSe-based Solar Cells [3](#cdtecdse-based-solar-cells)

1.2 CuInSe~2~-based Solar Cells [4](#cuinse2-based-solar-cells)

1.3 Si-based Solar Cells [4](#si-based-solar-cells-1)

1.4 In-line Diagnostics [5](#in-line-diagnostics)

1.5 Training and Education [5](#training-and-education)

1.6 Publications [5](#publications)

1.7 Organization of the Report [6](#organization-of-the-report)

2 CdTe/CdS-BASED SOLAR CELLS [7](#cdtecds-based-solar-cells-1)

2.1 CdTe Surface Chemistry [7](#cdte-surface-chemistry)

2.1.1 Effects of CdTe Etching on Device Performance
[7](#effects-of-cdte-etching-on-device-performance)

2.1.2 Formation of Copper Tellurides on Etched CdTe Surfaces
[10](#formation-of-copper-tellurides-on-etched-cdte-surfaces)

2.2 Contacts and Stability [12](#contacts-and-stability)

2.2.1 Systematic investigation of effects of humidity on CdTe devices
and components
[12](#systematic-investigation-of-effects-of-humidity-on-cdte-devices-and-components)

2.2.2 Effects of Humidity during Post-Deposition Processing
[18](#effects-of-humidity-during-post-deposition-processing)

2.2.3 Transparent Contacts [21](#transparent-contacts)

2.2.4 High throughput VT CdTe characterization and deposition studies
[22](#high-throughput-vt-cdte-characterization-and-deposition-studies)

2.2.5 Device results and process studies of high throughput VT CdTe
[24](#device-results-and-process-studies-of-high-throughput-vt-cdte)

2.4 Improving CdTe/CdS Solar Cell Performance
[29](#improving-cdtecds-solar-cell-performance)

2.4.1 Alternative CdS Deposition [29](#alternative-cds-deposition)

2.4.2 Optimization of CSD CdS [37](#optimization-of-csd-cds)

2.4.3 Microscale Junction Analysis [42](#microscale-junction-analysis)

2.4.4 CdTe Conductivity and Alternative Treatments
[45](#cdte-conductivity-and-alternative-treatments)

2.4.5 Effects of Cu as an impurity in CdS
[47](#effects-of-cu-as-an-impurity-in-cds)

2.5 National CdTe Team Activity and Collaborations
[52](#national-cdte-team-activity-and-collaborations)

3 CuInSe~2~-BASED SOLAR CELLS [53](#cuinse2-based-solar-cells-1)

3.1 In-line deposition Introduction
[53](#in-line-deposition-introduction)

3.1.1 Cracking and Adhesion [54](#cracking-and-adhesion)

3.1.2 System Improvements [54](#system-improvements)

3.2 CuInSe~2~-based Devices with Increased V~oc~
[57](#cuinse2-based-devices-with-increased-voc)

3.2.1 Cu(InAl)Se~2~ Solar Cells [57](#cuinalse2-solar-cells)

3.2.2 Experimental Procedures for Cu(InAl)Se~2~
[58](#experimental-procedures-for-cuinalse2)

3.2.3 Results: Material Characterization
[59](#results-material-characterization)

3.2.4 Results: Cu(InAl)Se~2~ Solar Cells
[62](#results-cuinalse2-solar-cells)

3.2.5 Cu(InAl)Se~2~ Device Characterization
[63](#cuinalse2-device-characterization)

3.2.5 Cu(InAl)Se~2~ Discussion And Conclusions
[65](#cuinalse2-discussion-and-conclusions)

3.2.6 CuAlSe~2~ Deposition and Characterization
[66](#cualse2-deposition-and-characterization)

3.2.7 Sulfur Diffusion in Cu(InGa)Se~2~
[67](#sulfur-diffusion-in-cuingase2)

3.3. Fundamental Interface and Material Characterization
[72](#fundamental-interface-and-material-characterization)

3.3.1 Emitter Layer Deposition and Junction Characterization
[72](#emitter-layer-deposition-and-junction-characterization)

3.3.2 Alternative CdS Deposition [76](#alternative-cds-deposition-1)

3.3.3 Cu(InGa)Se~2~ Optical Characterization
[77](#cuingase2-optical-characterization)

3.3.4 Ellipsometry Measurement Strategy
[78](#ellipsometry-measurement-strategy)

3.3.5 Optical Modeling [79](#optical-modeling)

3.3.6 Cu(InGa)Se~2~ Ellipsometry: Results And Discussion
[84](#cuingase2-ellipsometry-results-and-discussion)

3.3.7 Cu(InGa)Se~2~ Ellipsometry: Conclusions
[94](#cuingase2-ellipsometry-conclusions)

3.3.8 Team Activities and Collaboration
[95](#team-activities-and-collaboration)

4 SI BASED SOLAR CELLS [97](#si-based-solar-cells-2)

4.1 Fabrication of Large Grain Si Films
[97](#fabrication-of-large-grain-si-films)

4.1.1 Grain size enhancement of HWCVD Si films at 430°C.
[97](#grain-size-enhancement-of-hwcvd-si-films-at-430c.)

4.1.2 Grain size enhancement of HWCVD Si films at 600°C.
[102](#grain-size-enhancement-of-hwcvd-si-films-at-600c.)

4.2 Solar Cell Development [104](#solar-cell-development)

4.2.1 Microcrystalline p-i-n solar cell by PECVD
[104](#microcrystalline-p-i-n-solar-cell-by-pecvd)

4.2.2 Solar Cells with µc-Si HWCVD or PECVD i-layers
[105](#solar-cells-with-µc-si-hwcvd-or-pecvd-i-layers)

4.2.3 Impurities in PECVD and HW films and devices
[108](#impurities-in-pecvd-and-hw-films-and-devices)

4.2.4 Boron doped HWCVD Si films [112](#boron-doped-hwcvd-si-films)

4.2.5 National Thin Si Team Activities
[115](#national-thin-si-team-activities)

5 DIAGNOSTICS [116](#diagnostics-1)

5.1. Contact Wetting Angle [116](#contact-wetting-angle)

5.2 Ellipsometry [120](#ellipsometry)

5.3 Glancing Incidence X-Ray Diffraction
[121](#glancing-incidence-x-ray-diffraction)

6 ABSTRACT [124](#abstract)

7 REFERENCES [127](#references)

# 

# **LIST OF FIGURES**

Figure 1. Plot of R~oc~, before and after stressing at 85°C in air and
in the dark, vs. etch time for devices that received NP etching during
back contact processing. The R~oc~ values, before and after stress, for
the device etched in 0.05% Br~2~/methanol are also shown.
[9](#__RefHeading___Toc469989132)

Figure 2. GIXRD plot, at 2**°** incidence, of a 60 sec NP etched CdTe
film obtained, following removal of a Cu powder/graphite paste contact.
[11](#__RefHeading___Toc469989133)

Figure 3. Schematic diagram of apparatus for humid stress experiment.
[14](#__RefHeading___Toc469989134)

Figure 4. GIXRD plots of Canrom CdTe following stress at 80°C in
[15](#__RefHeading___Toc469989135)

Figure 5. GIXRD plots of BDH etched IEC CdTe following
[15](#__RefHeading___Toc469989136)

Figure 6. GIXRD plots of non-annealed 50 nm Au/30 nm Cu/50 nm Te
[16](#__RefHeading___Toc469989137)

Figure 7. XRD scans of CdCl~2~ powder re-measured during exposure to
humid air. [19](#__RefHeading___Toc469989138)

Figure 8. GIXRD scans of VT CdTe surface after vapor CdCl~2~ treatment
[20](#__RefHeading___Toc469989139)

Figure 9. GIXRD scans of VT CdTe surface after vapor CdCl~2~ treatment
[20](#__RefHeading___Toc469989140)

Figure 10. GIXRD pattern of ZnTe:Cu/CdTe/CdS/TEC-15
[22](#__RefHeading___Toc469989141)

Figure 11. AFM image of VT86 [24](#__RefHeading___Toc469989142)

Figure 12. AFM image of VT87 [24](#__RefHeading___Toc469989143)

Figure 13. AFM image of VT88 [24](#__RefHeading___Toc469989144)

Figure 14. AFM of VT89 [24](#__RefHeading___Toc469989145)

Figure 15. Dependence of efficiency on the time of the CdCl~2~ treatment
[26](#__RefHeading___Toc469989146)

Figure 16. QE at 0V for devices from VT85 with 2 values CdCl~2~
[27](#__RefHeading___Toc469989147)

Figure 17. Light and dark J-V curve on VT85.33-5.
[28](#__RefHeading___Toc469989148)

Figure 18. Scatter plot of QE at 450 nm vs V~oc~ for 40 CdTe devices.
[29](#__RefHeading___Toc469989149)

Figure 19. Schematic view of CSD deposition of CdS films.
[30](#__RefHeading___Toc469989150)

Figure 20. Solution surface temperature versus time for
[33](#__RefHeading___Toc469989151)

Figure 21. Tapping AFM images of 20 nm thick CSD CdS films deposited
[34](#__RefHeading___Toc469989152)

Figure 22. GIXRD pattern of 50 nm thick CSD film
[35](#__RefHeading___Toc469989153)

Figure 23. Normalized optical transmission of Corning 7059 borosilicate
glass, [36](#__RefHeading___Toc469989154)

Figure 24. Normalized optical transmission, T/(1-R), of window layer
used in VT cells. [38](#__RefHeading___Toc469989155)

Figure 25. AFM scans of CSD CdS film: as-deposited (left)
[39](#__RefHeading___Toc469989156)

Figure 26. Effect of CdCl~2~ treatment time of CdS layers on
[40](#__RefHeading___Toc469989157)

Figure 27. QE of 4 pieces from VT 84 as described in Table 13.
[42](#__RefHeading___Toc469989158)

Figure 28. Schematic diagram (not to scale) of taper mounted
[43](#__RefHeading___Toc469989159)

Figure 29. Schematic of solar cell mounted in epoxy at an angle,
exposing the junction. [44](#__RefHeading___Toc469989160)

Figure 30. Tapping AFM image of the CdTe-CdS interface region
[44](#__RefHeading___Toc469989161)

Figure 31. AQE of IEC CdS/CdTe device at +0.6 V bias with 4 bias light
conditions [48](#__RefHeading___Toc469989162)

Figure 32. Ratio of AQE with red bias to AQE with white bias for various
CdS/CdTe devices plus one CdTe device without CdS. Data from First Solar
divided by 10. [48](#__RefHeading___Toc469989163)

Figure 33. Ratio of AQE near V~MP~ with red light to white light for
[49](#__RefHeading___Toc469989164)

Figure 34. AQE of IEC device with red bias at various chopping
frequency. [50](#__RefHeading___Toc469989165)

Figure 35. Ratio of AQE with blue bias to no bias for same devices as
Figure 32. [50](#__RefHeading___Toc469989166)

Figure 36. Cracking of Mo/CIGS structure on Upilex substrate;
[53](#__RefHeading___Toc469989167)

Figure 37. Source layout in the CIGS in-line deposition system.
[55](#__RefHeading___Toc469989168)

Figure 38. Temperature of the sources and of the Se branches during a
CIGS run where heaters were properly sequenced during heat-up and
cool-down to keep Se source temperature the lowest. Right hand scale
shows temperature [56](#__RefHeading___Toc469989169)

Figure 39. Differences in grain size and morphology with different
[60](#__RefHeading___Toc469989170)

Figure 40. Load needed to remove a stud epoxied to a
[62](#__RefHeading___Toc469989171)

Figure 41. Light and dark J-V curves with Cu(InAl)Se~2~
[64](#__RefHeading___Toc469989172)

Figure 42. QE curves of devices in Figure 38. Solid lines were
[64](#__RefHeading___Toc469989173)

Figure 43. Ratio of QE(-1V) / QE(0V) showing increased
[65](#__RefHeading___Toc469989174)

Figure 44. SEM images of CuInSe~2~ crystals after sulfur reaction.
[68](#__RefHeading___Toc469989175)

Figure 45. AES depth profiles of the relative S content y measured on
crystals sulfurized for 1 hour and 9 hours, respectively. The x-axis
shows the sputter depth divided by the square root of the sulfurization
time. The fact that both curves agree in this graph confirms that the S
distribution in the crystals was shaped by one-dimensional diffusion.
[70](#__RefHeading___Toc469989176)

Figure 46. D~l~ as a function of y in CuIn(Se~1-y~S~y~)~2~, calculated,
with the help of the Boltzmann-Matano method, from the AES depth profile
of the crystal sulfurized [71](#__RefHeading___Toc469989177)

Figure 47. Comparison of the AES depth profile measured on
[71](#__RefHeading___Toc469989178)

Figure 48. Normalized transmissions of CBD ZnS film with various air
[73](#__RefHeading___Toc469989179)

Figure 49. Bandgap determination from (αE)^2^ versus E for the
[74](#__RefHeading___Toc469989180)

Figure 50. GXRD patterns: 1) as deposited ZnS/Mo; 2) air annealed
[75](#__RefHeading___Toc469989181)

Figure 51. Sample structure and optical model
[81](#__RefHeading___Toc469989182)

Figure 52. Schematic diagram showing how the crystal field
[82](#__RefHeading___Toc469989183)

Figure 53. The ε~2~ spectrum for the CuInSe~2~ thin film along with the
[86](#__RefHeading___Toc469989184)

Figure 54. Initial spectrum (dashed line) for a typical
CuIn~1-x~Ga~x~Se~2~ [87](#__RefHeading___Toc469989185)

Figure 55. Measurement and fit of the ellipsometry parameters
[88](#__RefHeading___Toc469989186)

Figure 56. Second derivatives of the ε~2~ spectrum for a CuInSe~2~ film
[90](#__RefHeading___Toc469989187)

Figure 57. Index of refraction (a) and extinction coefficient
[90](#__RefHeading___Toc469989188)

Figure 58. Second derivative of the ε~2~ spectra for CuIn~1-x~Ga~x~Se~2~
films [91](#__RefHeading___Toc469989189)

Figure 59. Variation in the fundamental absorption transitions
[91](#__RefHeading___Toc469989190)

Figure 60. Variation in the ∆~cf~ and ∆~so~ of CuIn~1-x~Ga~x~Se~2~ films
with x. [95](#__RefHeading___Toc469989191)

Figure 61. Variation in the higher energy transitions of
CuIn~1-x~Ga~x~Se~2~ films with x. [96](#__RefHeading___Toc469989192)

Figure 62. XRD pattern of poly-Si on glass, and poly-Si on Al coated
glass. [98](#__RefHeading___Toc469989193)

Figure 63. Comparison of Raman spectra of c-Si, poly-Si on glass,
[99](#__RefHeading___Toc469989194)

Figure 64. AFM image of the surface of a 2 μm thick poly-Si film
deposited Al. [100](#__RefHeading___Toc469989195)

Figure 65. XTEM image of 2 μm thick poly-Si film on Al coated glass.
[101](#__RefHeading___Toc469989196)

Figure 66. SEM planar image after etching Si film deposited on Al.
[101](#__RefHeading___Toc469989197)

Figure 67. Optical micrograph of back surface of Si film deposited on
Al. [101](#__RefHeading___Toc469989198)

Figure 68. Planar SEM/EDS of Si film deposited on Al at 600°C.
[103](#__RefHeading___Toc469989199)

Figure 69. Raman profiles of Al rich regions in films deposited at
600°C. [103](#__RefHeading___Toc469989200)

Figure 70. Log QE vs energy for a-Si, µc-Si and c-Si devices.
[104](#__RefHeading___Toc469989201)

Figure 71. Light and dark JV curves for n+-p+ cells GD41015-21
[107](#__RefHeading___Toc469989202)

Figure 72. QE at 0V, dark for 4 devices.
[108](#__RefHeading___Toc469989203)

Figure 73. QE ratios comparing forward bias to 0V on two devices
[109](#__RefHeading___Toc469989204)

Figure 74. Log QE plots showing increased midgap
[109](#__RefHeading___Toc469989205)

Figure 75. XRD pattern of B-doped HWCVD μc-Si film.
[113](#__RefHeading___Toc469989206)

Figure 76. Raman spectra of B-doped HWCVD μc-Si film.
[113](#__RefHeading___Toc469989207)

Figure 77. Deionized water droplet on the surface of VT CdTe.
[118](#__RefHeading___Toc469989208)

**LIST OF TABLES**

Table 1. Solvents used to monitor effect of changing Br~2~ etch solution
viscosity [7](#__RefHeading___Toc469989095)

Table 2. Sample set for humid stress experiment
[13](#__RefHeading___Toc469989096)

Table 3. Humidity conditions for each control agent at 80°C.
[13](#__RefHeading___Toc469989097)

Table 4. J-V data for Au/Cu/Te contacted IEC devices with CdTe made by
vapor transport, before and after stress at 80°C in various humidities
for 500 hrs. [17](#__RefHeading___Toc469989098)

Table 5. J-V data for First Solar devices before and after stress at
80°C in various humidities for 500 hrs.
[18](#__RefHeading___Toc469989099)

Table 6. Key deposition parameters used in depositing VT CdTe films
[23](#__RefHeading___Toc469989100)

Table 7. Effect of gas flow on thickness and growth rate of VT CdTe
films. All other parameters were as listed for baseline films in Table
6. VT86 was deposited at baseline conditions.
[23](#__RefHeading___Toc469989101)

Table 8. Efficiency and yield of pieces with different spacing of
CdCl~2~ plate and [27](#__RefHeading___Toc469989102)

Table 9. Time to complete reaction. [31](#__RefHeading___Toc469989103)

Table 10. CdS thickness (nm) for different substrate temperatures and
Cd^++^ concentration. [32](#__RefHeading___Toc469989104)

Table 11. Utilization of Cd^++^ (%) [33](#__RefHeading___Toc469989105)

Table 12. CdTe/CdS device J-V results.
[37](#__RefHeading___Toc469989106)

Table 13. Performance of best cell-on-pieces from VT 84 having
differently treated CdS. Samples CdS from FS included. Yield based on
shunting and low FF of 8 cells per piece.
[41](#__RefHeading___Toc469989107)

Table 14. Conductivity results for treated CdTe films.
[46](#__RefHeading___Toc469989108)

Table 15. Cu(InAl)Se~2~ device parameters for films with E~g~ = 1.1 to
1.5 deposited with different processes. With the 3-step process films
with E~g~ \> 1.1 did not adhere during device processing.
[63](#__RefHeading___Toc469989109)

Table 16. Energy Bandgap of CBD ZnS and sputtered ZnO
[73](#__RefHeading___Toc469989110)

Table 17. Summary of device results with different annealing times for
Cu(InGa)Se~2~ cells with CBD ZnS. Results include the best cell and
average of 6 cells on each piece. [76](#__RefHeading___Toc469989111)

Table 18. Cu(InGa)Se~2~/CdS device results using chemical surface
deposition for the CdS. [77](#__RefHeading___Toc469989112)

Table 19. EDS data of Cu(In~1-x~Ga~x~)Se~2~ films used for the SE
analysis. [85](#__RefHeading___Toc469989113)

Table 20. Comparison of orientation parameter, roughness data from
[85](#__RefHeading___Toc469989114)

Table 21. Oscillator energies corresponds to the electronic transitions
with varying Ga/(Ga+In) ratio. [88](#__RefHeading___Toc469989115)

Table 22. Optical constants *n* and *k* for Cu(InGa)Se~2~ with varying
x. [92](#__RefHeading___Toc469989116)

Table 23. Variation in the electronic transition energies with
Ga/(Ga+In) ratio obtained from Figure 59.
[93](#__RefHeading___Toc469989117)

Table 24. Properties of films deposited at 430°C.
[100](#__RefHeading___Toc469989118)

Table 25. Properties of films Deposited at 600°C.
[102](#__RefHeading___Toc469989119)

Table 26. Si i-layer deposition conditions.
[105](#__RefHeading___Toc469989120)

Table 27. Undoped a-Si buffer and Si p-layer deposition conditions.
[106](#__RefHeading___Toc469989121)

Table 28. Device structures and JV data.
[106](#__RefHeading___Toc469989122)

Table 29. JV results from n+-p+ cells having no GD or HW absorber.
[106](#__RefHeading___Toc469989123)

Table 30. Summary of SIMS for IEC Si layers and devices.
[111](#__RefHeading___Toc469989124)

Table 31. B-doped Si film deposition conditions.
[112](#__RefHeading___Toc469989125)

Table 32. Conductivity of B-doped Si films deposited at 200 mT.
[114](#__RefHeading___Toc469989126)

Table 33. Deposition conditions and film properties of B-doped Si films.
[115](#__RefHeading___Toc469989127)

Table 34. Properties of contact liquids
[117](#__RefHeading___Toc469989128)

Table 35. Contact wetting angle and surface energy for Cu(InGa)Se~2~
films. [119](#__RefHeading___Toc469989129)

Table 36. GIXRD and contact angle results for CdTe surfaces at different
stages of post-deposition processing. Samples marked with (\*) are the
same specimen. [120](#__RefHeading___Toc469989130)

Table 37. Mass absorption coefficient (μ/ρ) and sampling depth (τ) for
1° incident beam angle and different x-ray sources for thin-film
photovoltaic absorber layer materials.
[123](#__RefHeading___Toc469989131)

#  **1 Objectives**

The primary objectives of this Institute of Energy Conversion (IEC)
research are specific to the 3 types of thin film solar cells currently
under commercial interest. In addition, generic process diagnostic and
control tools will be developed.

## 1.1 CdTe/CdSe-based Solar Cells

The CdTe effort will address increased voltage and stability, the
development of an in-line CdTe process to implement high rate and high
throughput deposition on a moving substrate, and fundamental studies of
CdTe surface chemistry. To meet these objectives, IEC will:

-   Separate and quantify the effect on V~OC~ of high resistance (HR)
    buffer layers such as Ga~2~O~3~ and In~2~O~3~ and widening the CdS
    window layer bandgap with Zn.

-   Develop and apply a transparent back contact to CdTe to probe back
    contact junction formation and degradation.

-   Evaluate methods for increasing V~OC~ by reducing space charge
    recombination.

-   Improve the stability of CdTe devices by developing alternative
    processing steps to 'stabilize' the junction and contact; and to
    compare the effects of well-controlled stress conditions on CdTe
    solar cells by performing detailed device analysis and relating
    results to differences in fabrication.

-   Develop the fundamental understanding needed for high throughput
    CdTe module fabrication by controlling high rate delivery of Cd and
    Te species to the moving substrate surface.

-   Develop rapid and reproducible post-deposition processing
    alternatives to the current CdCl~2~ vapor treatment.

-   Determine fundamental understanding of etching and wet processing
    steps on CdTe surfaces

-   Determine relation between the contact processing and CdTe surface
    chemistry, and the device performance and stability especially as
    applies to First Solar devices.

## 1.2 CuInSe~2~-based Solar Cells

The CuInSe~2~-based effort will characterize fundamental properties of
interfaces in CuInSe~2~-based solar cells, develop approaches to
improving V~OC~ including alloying with S and Al, and will also include
the development of an in-line process for deposition on a moving
substrate. To meet these objectives, IEC will:

-   Develop the in-line deposition of Cu(InGa)Se~2~ films from
    stationary elemental evaporation sources onto moving rigid (Mo
    coated glass) and flexible (Mo coated polyimide) substrates at rates
    commensurate with commercial manufacturing.

-   Determine the relationship of the film properties, both
    compositional and structural, to the solar cell performance with
    particular emphasis on the effect of complex band-gap profiles due
    to the depth distribution of Ga and In and incorporation of Na.

-   Increase the operating voltage by increasing the bandgap, E~g~, in
    the absorber layer or in the space charge region using Cu(InAl)Se~2~
    alloys.

-   Investigate the feasibility of controlled p--type doping of the
    Cu(InGa)Se~2~ or Cu(InAl)Se~2~ absorber layers to increase V~OC~.

-   Provide a fundamental understanding of surface reactions and
    interface chemistry in the fabrication of CIGS devices using atomic
    level surface characterization techniques to study the Mo surface,
    Mo/ Cu(InGa)Se~2~ interface, the free Cu(InGa)Se~2~ surface and
    Cu(InGa)Se~2~/CdS interface.

-   Determine reaction pathways for the formation of Cu(InGa)Se~2~ and
    Cu(InGa)(SeS)~2~ on H~2~Se/H~2~S time-temperature-gas concentration
    profiles using sputtered Cu/Ga/In precursors.

-   Fabricate devices with CdS, ZnS, and CdZnS buffer layers and
    CuInSe~2~ with Ga, S, and/or Al alloy absorber layers in order to
    characterize the effect buffer layers on device behavior and
    fundamental interface properties.

## 1.3 Si-based Solar Cells

The Si-based effort will focus on developing a process for fabricating
polycrystalline Si solar cells and materials at low temperatures on low
cost substrates as well as the use of microcrystalline Si layers in
solar cells. To meet these objectives, IEC will:

-   Develop in-situ and/or post-deposition methods such as metal induced
    crystallization (MIC) to produce HWCVD deposited Si films and
    devices with grains in the micrometer range.

-   Investigate various forms of Eutectic Promoted Deposition (EPD)
    where very thin metal layers deposited on glass substrates are
    briefly heated above the metal-Si eutectic as the HWCVD Si
    deposition proceeds leading to precipitation of Si from the
    supersaturated liquid solution.

-   Investigate the feasibility of layer-by-layer deposition (LBL) to
    create thin Si seed layers in-situ to enhance nucleation of larger
    grains during subsequent growth.

-   Develop all HWCVD p n junction devices by developing HWCVD emitter
    layers on large grain HWCVD absorbers.

-   Complete work already underway of incorporating two-phase µc-Si(H,
    O, C) PECVD p-layers into superstrate p-i-n solar cells with a-Si
    i-layers to increase V~oc~ and blue response.

## 1.4 In-line Diagnostics

The in-line process diagnostic effort will develop diagnostic tools
needed for process control and the associated quantitative models that
link sensor outputs to process variables and material properties. To
meet these objectives, IEC will:

-   Understand the fundamentals of thin film growth and identification
    of critical properties that lead to efficient solar cells, i.e.,
    product specifications.

-   Identifying process parameters, to which film properties are
    sensitive, i.e., process determinants.

-   Develop robust, reliable and fast-response sensor equipment for
    in-situ applications in deposition reactors.

-   Develop quantitative models that relate sensor output to process
    determinants and film properties for use in model based process
    control, i.e., intelligent process control.

## 1.5 Training and Education

During the period of this subcontract (August 24, 1902 to August 23,
2003) IEC provided training and education for the following: one
visiting professionals; nine post-doctoral fellows; six graduate
students; and nine undergraduate students. One of our visiting
professionals was a Fulbright scholar. Names are given in the list of
contributors.

## 1.6 Publications

As a result of research performed under this contract, IEC published 21
papers.

## 

## 1.7 Organization of the Report

This report is organized into four technical sections: CdTe-based solar
cells, CuInSe~2~-based solar cells, Si-based solar cells, and
diagnostics. Each section describes the progress made at IEC in
addressing the critical issues discussed above during the 12-months
period of this contract.

# **2 CdTe/CdS-BASED SOLAR CELLS**

## 2.1 CdTe Surface Chemistry

### 2.1.1 Effects of CdTe Etching on Device Performance

Investigations of the dynamics of CdTe etching continued. Currently we
are attempting to systematically obtain accurate thicknesses, using
VASE, of Te films produced by NP etching. VASE measurements and modeling
are complicated by surface roughness effects, which are enhanced by NP
grain boundary (GB) etching. To minimize these effects the CdTe samples
have been chemically polished, by etching in 0.05% bromine methanol for
60 sec. VASE has shown BM etching to be very consistent and
reproducible, giving very similar surfaces on CdTe films from different
batches and of different ages. This is a very important result, which
allows confidence in the nature of the surface prior to the NP etch. The
presence of the BM generated thin Te-rich surface does not appear to
affect the reactivity of the NP etch with CdTe. The pieces described
above will also be systematically analyzed using other surface sensitive
techniques, and will be compared to unpolished samples, to further
understand the chemistry of CdTe and Te during NP etching.

In the previous annual report [^1], a new set of experiments was
described where the effects of changing the viscosity of Br~2~-based
etches on CdTe morphology were investigated. Increasing etch solution
viscosity was monitored in order to examine the potential for developing
wet-chemical approaches to produce Te-rich CdTe surfaces without
significant GB etching. Br~2~ etches were prepared with a series of
solvents of different viscosities (see Table 1). It was shown that CdTe
surfaces were smoothed following etching for 10 sec in each of the ∼0.1%
Br~2~/solvent solutions. As expected, surface roughness decreased with
decreasing viscosity. All solutions etched the CdTe surface to some
degree, however little difference in surface roughness was observed
following etching in solutions of viscosity \> ∼20 mPa s.

[]{#__RefHeading___Toc469989095 .anchor}Table 1. Solvents used to
monitor effect of changing Br~2~ etch solution viscosity

  ------------------- ------------------------- -------------------------
  Solvent             Viscosity (25°C) /mPa s   Device No.

  methanol            0.5                       VT58-321

  1-butanol           2.5                       VT58-322

  2-ethyl-1-hexanol   6.3                       VT58-331

  1-decanol           10.9                      VT58-432

  ethylene glycol     16.1                      VT58-332

  diethylene glycol   30.2                      VT58-422

  propylene glycol    40.4                      VT58-421

  glycerol            930                       VT58-431
  ------------------- ------------------------- -------------------------

A series of CdTe devices were processed with etching for 20 sec in
∼0.05% Br~2~ solutions of each of the solvents in Table 1. Previous
experiments investigating the effect of different HNO~3~/H~3~PO~4~ (NP)
etches on device performance, found that for treatments which produce
low amounts of surface-Te, preparing the back contact by deposition of
10 nm Cu films followed by application of graphite paste resulted in
significant device shunting, most likely due to an excess of Cu. To
avoid this problem for these Br~2~-etched devices, contacts were
prepared using graphite paste containing Cu-powder. It is postulated
that the Te layer would react only with stoichiometric levels of Cu and
diffusion of excess Cu from the contact paste into the cell structure
would be slowed.

The devices were contacted with graphite paste containing ∼6% w/w Cu
powder and annealed at 200°C in Ar(g) for 30 min. Immediately following
processing, all devices were alive and showed similar J‑V behavior. J-V
characteristics were, however, generally poor, e.g. V~oc~∼700 mV,
J~sc~∼15 mA cm^‑2^ and efficiency∼6.5%. This is not unexpected as back
contact processing was not optimized for the etch treatments. We also
note V~oc~∼700 mV has been obtained for other CdTe devices contacted
with Cu-containing graphite paste in our laboratory. FF of the cells
were reasonable, ∼55-60%, though the FF of the device etched in
Br~2~/2-ethyl-1-hexanol was \<50% due to high R~oc~. Following resting
overnight, the devices etched in 1-decanol, diethylene glycol, propylene
glycol and glycerol solutions, i.e. generally the higher viscosity
solvents, were found to have become highly shunted or completely
shorted. The J-V characteristics of the other devices were unchanged.
This observation highlights the very fast GB diffusion dynamics of Cu in
CdTe films. High viscosity Br~2~ etches are expected to produce thinner
Te layers than the less viscous etches over the same treatment time. The
thinner films will react quickly with Cu but will not prevent or
significantly slow diffusion of excess Cu. Further experiments are
planned to optimize the Br~2~ etches, including etch time and solution
concentrations at different viscosities. Effects of longer etch times on
GB integrity and device performance will also be monitored.

For comparison, a series of CdTe/CdS devices were processed following
etching in NP for 20, 30, 45, 60 sec and multiple etches of 3x20s and
2x30s and one device received a 20 sec ∼0.05% BM etch. Devices were
completed with graphite pastes containing ∼6% w/w Cu powder and the
contacts were annealed for 30 min at 200°C in an Ar(g) atmosphere. All
the completed devices again showed similar J-V characteristics V~oc~
∼700 mV, J~sc~ ∼ 15 mA cm^-2^, fill factor ∼55-60% and efficiency ∼
6-6.5%, which again are consistent with other devices prepared with this
paste. R~oc~ was also similar for each of the single-NP and -BM etched
devices (Figure 1), and was slightly higher for devices receiving
shorter NP etch. The BM etched device had a R~oc~ similar to the devices
etched for 45 and 60 sec in NP.

These devices were stressed in air at 85°C in the dark and were removed
for J-V testing every 7 -- 10 days. Monitoring of the stressed devices
during the first 30-40 days showed that very little degradation occurred
in any of the cells. Following 50 days stress, some degradation was
observed in the devices that had received shorter NP etches. The
degradation was the result of an increase in R~oc~ (Figure 1), as no
V~oc~ or J~sc~ losses, or rollover were observed for any piece. The
device which received a 60 sec NP etch showed only a very slight
increase in R~oc~, while the BM etched device exhibited a significant
increase in R~oc~, though less than that observed for the 20 sec
NP-etched sample.

![](media/image2.png){width="3.2708661417322835in"
height="3.239763779527559in"}

[]{#__RefHeading___Toc469989132 .anchor}Figure 1. Plot of R~oc~, before
and after stressing at 85°C in air and in the dark, vs. etch time for
devices that received NP etching during back contact processing. The
R~oc~ values, before and after stress, for the device etched in 0.05%
Br~2~/methanol are also shown.

The cells that received multiple etches in NP exhibited similar initial
J-V behavior to the single etched batch, though the 2x30 sec etched
device was slightly superior. As with the single etched devices, little
degradation was observed following stress, except for an increase in
R~oc~. Interestingly, the device which received 3x20 sec NP etches
showed a significant increase in R~oc~, similar to the device which
received a single 20 sec etch, while the 2x30 sec etched cell showed an
R~oc~ increase similar to the 30 sec and 45 sec etched pieces. In the
previous annual report, it was shown that multiple NP etches of 20 sec,
with the film being removed prior to the signature bubbling at 20-25 sec
immersion, showed only small additional Te formation with each
subsequent etch, with the Te thickness plateauing following 6-7
immersions. At this point the amount of Te generated was still
significantly less than that produced by a single 30 sec NP etch.
Subsequent 30 sec etches continue to produce more Te following 2-3
immersions, before the Te begins to be consumed by the solution with
further etching. Therefore, we can assume that the 3x20 sec and 2x30 sec
NP etched cells had Te-rich surface films similar in thickness to the 20
sec and 45 sec NP etched devices, respectively. This was qualitatively
confirmed by GIXRD.

The action of NP etches on CdTe is known to produce Te-rich films of
∼50-100 nm thickness, along with significant etching and widening of the
GBs. By comparison, BM etching produces Te-rich films of only a few
angstroms in thickness. A longer NP etch is expected to produce a
thicker Te film, but also a higher degree of GB etching. Extensive
etching of GBs may influence the diffusivity of Cu and other impurities
throughout the device. However, the observed device results indicate
that the mild cell degradation observed in this study was due to back
contact effects, i.e. R~OC~ increase, and only occurred for devices
within thinner Te-rich layers or that received shorter etches. We
suggest that completed cells processed with back contacts containing
thicker Te layers may improve contact and device stability. Thicker Te
layers, coupled with the affinity of Te for Cu, would attenuate the
escape of Cu from the Cu~2-x~Te back contact region. That is, Te may
getter Cu and slows its release into the CdTe film. Eventually
significant Cu would be expected to diffuse from the Te layer and
diffuse through the cell, however, the overall rate of Cu diffusion
would be significantly slowed and thus maintain a 'copper telluride'
back contact for a longer period of time. The phase of 'copper
telluride' is not clear, however we suggest that a Cu-poor 'copper
telluride' phase is likely to be a more favorable back contact than pure
Te. As also discussed Section 2.2, thicker Te contacts are also likely
to compensate oxidation and other chemical changes occurring at the back
contact during stress.

### 2.1.2 Formation of Copper Tellurides on Etched CdTe Surfaces

The formation of copper tellurides on wet chemical etched CdTe surfaces
during back contacting has been monitored. Similar experiments have been
carried out at IEC in the past; however, these measurements were carried
out using deposited Cu and Te films of stoichiometric thicknesses. In
this work, copper telluride contacts were prepared by coating etched
CdTe surfaces with Cu-containing graphite pastes. GIXRD was used to
monitor the resultant surface products.

The CdTe surfaces were etched with either NP or 0.1% BM. The etched
surfaces were coated with graphite paste containing \~6% w/w Cu, CuCl or
CuI powders. The contacts were annealed at 200°C in Ar(g) for 30 min.
For GIXRD measurements, the graphite films were removed by sonication in
acetone. A purple-tinge was noted on the CdTe surface of some samples,
from both NP and BM etched, following removal of the graphite contact.

GIXRD confirmed that etching of CdTe in BM for 20 sec or in NP for 60
sec, removes oxides from the film surface. Both etches are known to
produce Te-rich surface films. Crystalline Te is detected with GIXRD
following NP etching. BM etching, however, produces a very thin and
amorphous Te-rich surface and no Te is observed with GIXRD.

Figure 2 shows the 22 -- 34.5° region of the GIXRD plot, at 2° incident
angle, of a 60 sec NP etched CdTe film following removal of the Cu
powder/graphite paste back contact. The CdTe 111 and Te 101 peaks are
present at 23.7° and 27.6°, respectively. A series of new peaks have
emerged at 25.3°, 26.7°, 31.2° and 33.9°. GIXRD of CdTe samples
contacted with Cu powder/graphite paste following 20 sec BM etching (not
shown), though lacking the Te reflection, showed similar peaks. However,
the peaks were very weak in comparison to the NP etched sample, except
for the 26.7° reflection, which was of similar intensity to that peak in
Figure 2. The 25.3° reflection is also now a doublet with a slightly
more intense peak at 25.7°.

![](media/image3.png){width="4.083333333333333in"
height="2.736111111111111in"}

[]{#__RefHeading___Toc469989133 .anchor}Figure 2. GIXRD plot, at 2°
incidence, of a 60 sec NP etched CdTe film obtained, following removal
of a Cu powder/graphite paste contact.

The peaks at 25.3°, 31.2° and 33.9° can be assigned to the formation of
a CuTe (JPCDS 22-0252) phase on both the NP and BM etch surfaces. The
25.7° reflection from the BM etched sample may also be due to CuTe. This
phase is clearly the dominant product on the NP etched surface and
contradicts previous work of back-contacts prepared from stoichiometric
Te and Cu films where formation of the Cu~2-x~Te phase was observed.
This is not unexpected due to the greater amount of Te generated by NP
etching coupled with the slow release of Cu from the graphite paste. The
weaker intensity of the reflections on the BM etched surface may be
expected due to the thinner Te-rich surface film. The peak at 26.7°
could be assigned to a Cu~3~Te~2~-related phase, however, the relative
intensity of the peak was found to be inconsistent between different
samples and the intensity was found to decrease with further sonication
of films. This could indicate that the Cu~3~Te~2~ phase, if present, may
not adhere well to the CdTe surface, however, the GIXRD spectrum of a
graphite paste film, dried on glass, shows graphite to have a single
strong reflection centered at 26.6°. Therefore, we can assume that this
peak in the GIXRD plots of the CdTe films is most likely due to residual
graphite remaining on the surface.

The residues of the carbon contact remaining after the sonication step
were found to be removed by wiping the surfaces of the CdTe films with a
Q-tip soaked in acetone. GIXRD measurements showed, for both BM and NP
etched samples, the graphite peak is removed leaving a single peak
centered at 26.95°. The assignment of this peak is unclear. A phase of
Cu~3~Te­~2~ is possible, as is a Cu~3~Te~4~ phase. Either of these phases
could account for the purple-tinge observed on the surface of some of
these samples

To investigate the reactivity of different forms of Cu in the back
contact paste towards CdTe surfaces, etched films were contacted with
graphite paste containing CuCl and CuI powders. On 20s BM etched CdTe,
both powders produced only very weak peaks at 26.6°, due to graphite,
and 27.6°, which is most likely due to the formation of a small amount
of Cu~2‑x~Te rather than crystalline Te. A 60 sec NP etched CdTe surface
contacted with CuCl-containing graphite paste showed peaks due to CuTe,
at 25.2° and 31.1°, crystalline Te and graphite. Again, CuTe is the
dominant phase produced on NP etched CdTe.

These results indicate that different copper telluride phases can be
produced during the back contacting process, depending on the chemical
state of Cu, the nature of application and the CdTe surface preparation.
These results also suggest that significant levels of Cu~2-x~Te may not
be required to produce quality back contacts to CdTe/CdS cells, and that
any copper telluride phase could produce a reasonable back contact.
However, it is likely that certain copper telluride phases may be
superior to others and produce better performing devices. The nature of
the Cu + Te reaction requires further investigation to confirm
conditions required for the formation of the various copper telluride
phases, which can then be correlated to different back contacting
processes and to device performance.

# 

## 2.2 Contacts and Stability

### 2.2.1 Systematic investigation of effects of humidity on CdTe devices and components

A number of films of device components and completed devices were
selected and stressed over a range of constant humidities at an elevated
temperature of 80°C. Table 2 shows a list of the samples selected for
this study. Films were selected from different sources and from
different modes of deposition. The Au/Cu/Te films duplicate the IEC
device back contacts, and consist of 50 nm Au/30 nm Cu/ 50 nm Te. The
IEC devices were etched with Br~2~-dichromate-hydrazine (BDH), giving
\~50 nm Te surface film, prior to deposition of the dot Au/Cu/Te
contacts (0.07 cm^2^ diameter). The contacted devices received a 30 min
anneal in Ar(g) at 190°C, while the Au/Cu/Te films were not annealed.
The First Solar cells are standard 'dot' devices with contacts
evaporated through a mask. Both types of devices were completed with In
solder at the front contact.

[]{#__RefHeading___Toc469989096 .anchor}Table 2. Sample set for humid
stress experiment

+------------------------+-------------------------+------------------+
| IEC                    | First Solar             | Canrom           |
+------------------------+-------------------------+------------------+
| Bare VT CdTe (CdCl~2~) | Bare FS VT CdTe         | Bare CdTe (no    |
|                        | (CdCl~2~)               | CdCl~2~)         |
+------------------------+-------------------------+------------------+
| Bare PVD CdTe          | BDH etched CdTe         |                  |
| (CdCl~2~)              |                         |                  |
+------------------------+-------------------------+------------------+
| Bare PVD CdS           | Bare CSS CdS            |                  |
+------------------------+-------------------------+------------------+
| Bare CSD CdS           | VT CdTe 'dot' cell (In  |                  |
|                        | solder)                 |                  |
+------------------------+-------------------------+------------------+
| Au/Cu/Te/glass (no     |                         |                  |
| anneal)                |                         |                  |
+------------------------+-------------------------+------------------+
| VT CdTe cell           |                         |                  |
|                        |                         |                  |
| (Au/Cu/Te contact, In  |                         |                  |
| solder)                |                         |                  |
+------------------------+-------------------------+------------------+

Humidity was controlled by storing the samples in a sealed vessel over
saturated solutions of various salts. This is a well-known method for
easy and inexpensive humidity control. Tables listing humidity with
temperature for different salt solutions are available. Table 3 shows
the salts and humidities, including saturation pressure and
concentration of H~2~O vapor in the vessel atmosphere. For the ∼0% and
100% humidities, the samples were stored over Drierite pieces (anhydrous
CaSO~4~) and pure H~2~O, respectively. Figure 3 shows the apparatus
set-up used. The samples were stored in a sealed tube with a check valve
to prevent over-pressure. Vials of the salt solutions were placed at the
base of the tube. Significant amounts of the respective salts were added
to the vials to ensure the solutions remained saturated throughout the
experiment. The ∼0% humidity samples were stored in a desiccator. The
conditions on loading were 22°C with 45% relative humidity. The tubes
and desiccator were placed in an oven and the check valves were closed
after 30 min. The vials were occasionally monitored to ensure that the
excess salt did not dissolve. Samples were stressed for 500 hrs. A
control set of samples was stored in ambient room air, ∼25°C at ∼30%
humidity, through this time.

[]{#__RefHeading___Toc469989097 .anchor}Table 3. Humidity conditions for
each control agent at 80°C.

+------------------+----------+-----+-----+-----+-----+--------------+
| Humidity Control | Drierite | S   | S   | S   | H~  | Room air     |
| Agent            |          | at. | at. | at. | 2~O | (control,    |
|                  |          |     |     |     |     | \~25°C)      |
|                  |          | L   | N   | KBr |     |              |
|                  |          | iCl | aBr |     |     |              |
+------------------+----------+-----+-----+-----+-----+--------------+
| Rel. humid%      | 0.0017   | 1   | 5   | 79  | 100 | ∼30          |
|                  |          | 0.1 | 1.4 |     |     |              |
+------------------+----------+-----+-----+-----+-----+--------------+
| P~sat~ (Torr)    | 0.006    | 37  | 182 | 280 | 355 |              |
+------------------+----------+-----+-----+-----+-----+--------------+
| \[H~2~O\] (g/L)  | 0.005    | 0   | 0   | 0   | 0   | ∼ 0.025      |
|                  | mg/L     | .03 | .15 | .23 | .29 |              |
+------------------+----------+-----+-----+-----+-----+--------------+

[]{#__RefHeading___Toc469989134 .anchor}Figure 3. Schematic diagram of
apparatus for humid stress experiment.

Following stressing, samples were removed and observed. Most pieces show
no visible changes, except for some CdTe pieces, which showed darkening
in drier atmospheres, and the Au/Cu/Te films and Au/Cu/Te contacts on
IEC devices, which became cloudy in the drier atmospheres.

Figure 4 shows the GIXRD patterns obtained from Canrom CdTe at different
humidities. These CdTe samples had not received CdCl~2~ treatment. From
the figure it can be seen that TeO~2~ had formed on the sample that had
been in the driest conditions, but no products are observed in the
wetter conditions or on the control piece. It would be expected that
surface species are very likely to form in the more humid systems,
therefore, we may assume that any products formed in these conditions
are likely to be amorphous Cd-Te oxide hydroxides or hydrates. Amorphous
species would not be detected with GIXRD. First Solar and IEC CdTe that
had received CdCl~2~ treatments (not shown) showed the expected presence
of ∼1500 Å CdO over all humidities, including the control. However,
significant levels of CdTeO~3~ were present only on the ∼0% and 10%
humidities and control pieces, suggesting that this phase becomes
hydrated and amorphous in the wetter conditions. The control BDH etched
IEC CdTe shows a crystalline Te peak (see Figure 5.) After stressing in
dry conditions, a large surface TeO~2~ component is observed. However,
the samples stressed in humidities of \>0% show no TeO~2~ and only very
weak diffuse Te peaks. This observation strongly suggests that amorphous
Te hydroxides or hydrates are formed during humid stressing.

> ![](media/image6.png){width="4.679861111111111in"
> height="3.441666666666667in"}

[]{#__RefHeading___Toc469989135 .anchor}Figure 4. GIXRD plots of Canrom
CdTe following stress at 80°C in

various humidities for 500 hrs.

> ![](media/image7.png){width="4.820138888888889in" height="3.54375in"}

[]{#__RefHeading___Toc469989136 .anchor}Figure 5. GIXRD plots of BDH
etched IEC CdTe following

stress at 80°C in various humidities for 500 hrs.

GIXRD analysis of the stressed IEC CSD CdS (not shown,) showed the
presence of CdO on all pieces, including the control, except for the
sample that had received stress at 100% humidity. The amount of CdO
detected generally decreased with increasing humidity stress levels.
These results suggest that CdO is formed on CSD CdS during growth, as
indicated by its presence on the control piece. However, with stress in
higher humidities, consistent with CdTe, it is suggested that the
surface oxides are converted to amorphous hydroxides or oxide hydrates.

Figure 6 shows the GIXRD patterns of the non-annealed Au/Cu/Te films
following humid stress. The control piece shows the presence of
Cu~2-x~Te, Cu~3~Te~2~ etc. and the Au film. Stressing at higher humidity
shows some changes in the composition of the copper telluride film,
which can be attributed to the films being slowly annealed at the
elevated temperature. At ∼0% humidity, however, the Au/Cu/Te film has
almost completely converted to TeO~2~, which results in the films
becoming cloudy, and highlights O~2~ corrosion of the copper telluride
back contact. This suggests that Cu~2-x~Te is either resistant to H~2~O
oxidation or that H~2~O oxidation dominates over O~2~ oxidation but is
self-limiting, corroding only the grain surfaces and not the entire
film.

> ![](media/image8.png){width="4.85in" height="3.561111111111111in"}

[]{#__RefHeading___Toc469989137 .anchor}Figure 6. GIXRD plots of
non-annealed 50 nm Au/30 nm Cu/50 nm Te

films following stress at 80°C in various humidities for 500 hrs.

GIXRD patterns of the completed IEC cells (not shown) showed similar
behavior to the Au/Cu/Te films. These devices received a BDH etch before
deposition of Au/Cu/Te dots. The GIXRD results showed that the Te film
between the contact dots also converted to TeO~2~ + Te at low humidity,
and probably amorphous hydrated phases at high humidity. This may
suggest that Te is slightly more resistant to O~2~ stress than humid
stress. Table 4 shows the J-V data for the IEC devices before and after
humid stress, including indications of up-down sweep retrace hysteresis,
extent of light-dark crossover (XOVR) and forward bias curvature.
Despite the extensive changes observed in the chemistry of the Au/Cu/Te
films and back contacts, the IEC devices show very stable electrical
behavior, with no significant changes in performance following stress at
all humidities. Only a slight, 5-20 mV, decrease in V~oc~ is observed.

[]{#__RefHeading___Toc469989098 .anchor}Table 4. J-V data for Au/Cu/Te
contacted IEC devices with CdTe made by vapor transport, before and
after stress at 80°C in various humidities for 500 hrs.

<table>
<colgroup>
<col style="width: 9%" />
<col style="width: 10%" />
<col style="width: 2%" />
<col style="width: 6%" />
<col style="width: 6%" />
<col style="width: 5%" />
<col style="width: 6%" />
<col style="width: 5%" />
<col style="width: 5%" />
<col style="width: 5%" />
<col style="width: 10%" />
<col style="width: 10%" />
<col style="width: 8%" />
<col style="width: 7%" />
</colgroup>
<tbody>
<tr class="odd">
<td>Sample</td>
<td>Condition</td>
<td>T</td>
<td>RH%</td>
<td>V<sub>oc</sub></td>
<td>J<sub>sc</sub></td>
<td>FF</td>
<td>Eff</td>
<td>R<sub>oc</sub></td>
<td>G<sub>sc</sub></td>
<td><p>Light</p>
<p>retrace</p>
<p>hysteresis</p></td>
<td><p>Dark</p>
<p>retrace</p>
<p>hysteresis</p></td>
<td><p>Light</p>
<p>dark</p>
<p>XOVR</p></td>
<td><p>Fwd</p>
<p>bias</p>
<p>Curv.</p></td>
</tr>
<tr class="even">
<td>VT84.43</td>
<td>Initial</td>
<td></td>
<td></td>
<td>743</td>
<td>20.5</td>
<td>67.0</td>
<td>10.2</td>
<td>4</td>
<td>2</td>
<td>N</td>
<td>N</td>
<td>N</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT84.43</td>
<td>Control</td>
<td>25</td>
<td>30</td>
<td>732</td>
<td>20.0</td>
<td>64.6</td>
<td>9.5</td>
<td>5</td>
<td>1</td>
<td>N</td>
<td>N</td>
<td>N</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT84.13</td>
<td>Initial</td>
<td></td>
<td></td>
<td>740</td>
<td>21.2</td>
<td>58.4</td>
<td>9.2</td>
<td>6</td>
<td>3</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT84.13</td>
<td>Drierite</td>
<td>85</td>
<td>0</td>
<td>736</td>
<td>22.7</td>
<td>57.1</td>
<td>9.6</td>
<td>6</td>
<td>3</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT84.33</td>
<td>Initial</td>
<td></td>
<td></td>
<td>748</td>
<td>22.4</td>
<td>65.1</td>
<td>11</td>
<td>4</td>
<td>1</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT84.33</td>
<td>LiCl</td>
<td>85</td>
<td>10.5</td>
<td>740</td>
<td>21.8</td>
<td>63.6</td>
<td>10.2</td>
<td>5</td>
<td>2</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT84.23</td>
<td>Initial</td>
<td></td>
<td></td>
<td>738</td>
<td>23.2</td>
<td>64.2</td>
<td>11</td>
<td>5</td>
<td>2</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT84.23</td>
<td>NaBr</td>
<td>85</td>
<td>51.4</td>
<td>725</td>
<td>22.3</td>
<td>60.6</td>
<td>9.7</td>
<td>6</td>
<td>2</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT84.32</td>
<td>Initial</td>
<td></td>
<td></td>
<td>742</td>
<td>21.9</td>
<td>61.9</td>
<td>10.1</td>
<td>5</td>
<td>3</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT84.32</td>
<td>KBr</td>
<td>85</td>
<td>79</td>
<td>721</td>
<td>19.6</td>
<td>60.2</td>
<td>8.5</td>
<td>6</td>
<td>3</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT84.22</td>
<td>Initial</td>
<td></td>
<td></td>
<td>702</td>
<td>22.6</td>
<td>59.5</td>
<td>9.4</td>
<td>6</td>
<td>2</td>
<td>N</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT84.22</td>
<td>H2O</td>
<td>85</td>
<td>100</td>
<td>699</td>
<td>19.2</td>
<td>56.6</td>
<td>7.6</td>
<td>7.0</td>
<td>2</td>
<td>N</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
</tr>
</tbody>
</table>

The only changes observed in the GIXRD results of First Solar devices
(not shown) was the formation of In~2~O~3~ on the In solder. At higher
humidities, GIXRD data indicated the In~­2~O~3~ became hydrated. No back
contact materials could be detected. J-V testing of the devices before
and after humid stress, showed significant changes in device
performance, with increasing R~oc~ and decreasing fill factor for all
stressed pieces and the control, a V~oc~ decrease of 20-40 mV for the
stressed pieces and forward bias curvature for the pieces stressed at
higher humidities (see Table 5). The change in device performance is not
expected to be due to the formation of In~2~O~3,~ as the probe would
have pierced the surface oxide layer during J-V testing.

[]{#__RefHeading___Toc469989099 .anchor}Table 5. J-V data for First
Solar devices before and after stress at 80°C in various humidities for
500 hrs.

<table>
<colgroup>
<col style="width: 8%" />
<col style="width: 10%" />
<col style="width: 3%" />
<col style="width: 6%" />
<col style="width: 5%" />
<col style="width: 5%" />
<col style="width: 5%" />
<col style="width: 4%" />
<col style="width: 5%" />
<col style="width: 5%" />
<col style="width: 10%" />
<col style="width: 9%" />
<col style="width: 9%" />
<col style="width: 9%" />
</colgroup>
<tbody>
<tr class="odd">
<td>Sample</td>
<td>Condition</td>
<td>T</td>
<td>RH%</td>
<td>V<sub>oc</sub></td>
<td>J<sub>sc</sub></td>
<td>FF</td>
<td>Eff</td>
<td>R<sub>oc</sub></td>
<td>G<sub>sc</sub></td>
<td><p>Light</p>
<p>retrace</p>
<p>hysteresis</p></td>
<td><p>Dark</p>
<p>retrace</p>
<p>hysteresis</p></td>
<td><p>Light</p>
<p>dark</p>
<p>XOVR</p></td>
<td><p>Fwd</p>
<p>bias</p>
<p>Curv.</p></td>
</tr>
<tr class="even">
<td>L2B.5</td>
<td>Initial</td>
<td></td>
<td></td>
<td>782</td>
<td>18.6</td>
<td>64.1</td>
<td>9.3</td>
<td>6</td>
<td>1</td>
<td>N</td>
<td>N</td>
<td>N</td>
<td>N</td>
</tr>
<tr class="odd">
<td>L2B.5</td>
<td>Control</td>
<td>25</td>
<td>30</td>
<td>821</td>
<td>19.0</td>
<td>58.3</td>
<td>9.1</td>
<td>11</td>
<td>1</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>L2B.4</td>
<td>Drierite</td>
<td>85</td>
<td>0</td>
<td>750</td>
<td>17.1</td>
<td>56.1</td>
<td>7.2</td>
<td>12</td>
<td>2</td>
<td>W</td>
<td>W</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>L2B.2</td>
<td>LiCl</td>
<td>85</td>
<td>10.5</td>
<td>775</td>
<td>17.6</td>
<td>56.5</td>
<td>7.7</td>
<td>12</td>
<td>1</td>
<td>W</td>
<td>W</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>L2B.3</td>
<td>NaBr</td>
<td>85</td>
<td>51.4</td>
<td>767</td>
<td>17.6</td>
<td>55.1</td>
<td>7.4</td>
<td>12</td>
<td>1</td>
<td>S</td>
<td>S</td>
<td>Y</td>
<td>Y</td>
</tr>
<tr class="odd">
<td>L2B.1</td>
<td>Initial</td>
<td></td>
<td></td>
<td>791</td>
<td>17.9</td>
<td>65.2</td>
<td>9.2</td>
<td>6</td>
<td>2</td>
<td>N</td>
<td>N</td>
<td>N</td>
<td></td>
</tr>
<tr class="even">
<td>L2B.1</td>
<td>KBr</td>
<td>85</td>
<td>79</td>
<td>755</td>
<td>17.1</td>
<td>54.2</td>
<td>7</td>
<td>14</td>
<td>2</td>
<td>S</td>
<td>S</td>
<td>Y</td>
<td>Y</td>
</tr>
<tr class="odd">
<td>L2B.6</td>
<td>H<sub>2</sub>O</td>
<td>85</td>
<td>100</td>
<td>770</td>
<td>17.6</td>
<td>57.7</td>
<td>7.8</td>
<td>11</td>
<td>1</td>
<td>S</td>
<td>S</td>
<td>Y</td>
<td>Y</td>
</tr>
</tbody>
</table>

The FS devices appear less robust than the IEC devices, which showed no
significant decrease in performance despite the extensive chemistry
occurring in the back contact. The IEC contact contains a very large
proportion, ∼100 nm, of Te, formed by a mixture of wet chemical etching
and evaporation. This thick Te component may allow the device to
compensate any chemical change in the contact, including oxidation
and/or Cu diffusion, and maintain a favorable working contact. Likewise,
we can speculate that a thick Te component would getter Cu at the back
contact for a longer period of time, thus increasing the contact
lifetime (see also Section 2.1. discussion). In contrast, no Te could be
detected by GIXRD in the First Solar back contacts. A very thin Te
component would be oxidized more easily and the contact would not
adequately compensate chemical or Cu-diffusion related changes. Changes
in the contact chemistry can be observed from the First Solar device
results.

These experiments are continuing. All stressed pieces are yet to be
analyzed by GIXRD. Subsequent techniques will be explored, e.g. XPS,
FTIR, SIMS etc., to attempt to detect the presence of amorphous
hydroxides and hydrates. A set of samples is currently being produced to
investigate the effect of humid stress testing under illumination.

### 2.2.2 Effects of Humidity during Post-Deposition Processing

### 

For post-deposition processing of cells using VT CdTe and CSD CdS, the
ambient condition during vapor CdCl~2~ treatment exerts a strong effect
on the surface chemistry, which is expected to affect bulk electronic
properties of the CdTe film. In particular, humidity in the treatment
ambient leads to conversion of the film surface to the native oxides
CdTeO~3~ and CdTe~2~O~5~. Elevated humidity levels can arise from the
ambient gas (e.g., room air) and from the CdCl~2~ source itself, which
is highly hygroscopic. Figure 7 is a sequence of XRD patterns acquired
from coarse, 100 m, CdCl~2~ powder (Alfa/Aesar 99.998% purity anhydrous
CdCl~2~) measured consecutively in room air. Each scan required 10
minutes for acquisition over the 2-theta range from 5 to 25 degrees.
Peaks due to cadmium chloride hydrate are detected in the fresh
specimen, which was exposed to air for only 10 minutes (mounting time
plus 5 minutes scan time). The hydrate peak intensity increased
approximately linearly with time; the peaks shift to higher angle due to
expansion of the powder out of the diffractometer focal plane. For the
480 min measurement, the powder had reacted into a different, unknown,
phase.

![](media/image9.png){width="5.116666666666666in"
height="3.251388888888889in"}

[]{#__RefHeading___Toc469989138 .anchor}Figure 7. XRD scans of CdCl~2~
powder re-measured during exposure to humid air.

During vapor CdCl~2~ treatment, dehydration of the source, water
desorption from the sample, and moisture content of the ambient all
contributes to the total quantity of water available for reaction at the
CdTe surface. GIXRD measurements of VT CdTe films made by IEC and First
Solar and vapor treated at IEC show that for vapor CdCl~2~ treatment in
dry ambient, with \[H~2~O\] \< 0.001 mg/l, only CdO is detected on the
surface. However, for treatments in ambient with \[H~2~O\] greater than
10 mg/l, native oxides of CdTe are formed, and little or no CdO is
formed. This is shown in Figures 8 and 9 below. The GIXRD scans for VT
films from a single deposition treated for 15, 30 and 60 minutes at
430°C in CdCl~2~ vapor at two different humidity levels. The strongest
peaks are due to CdTe. In Figure 8, the peaks at 33° and 55° are CdO.
The remaining peaks are the native oxide CdTeO~3~. In Figure 9**,** no
CdO is detected and all small peaks are due to CdTeO~3~ and CdTe~2~O~5~.

![](media/image10.png){width="5.173611111111111in"
height="2.734027777777778in"}

[]{#__RefHeading___Toc469989139 .anchor}Figure 8. GIXRD scans of VT CdTe
surface after vapor CdCl~2~ treatment

at 430°C for the times shown in air ambient with \~10 mg/l moisture
content.

![](media/image11.png){width="5.173611111111111in"
height="2.734027777777778in"}

[]{#__RefHeading___Toc469989140 .anchor}Figure 9. GIXRD scans of VT CdTe
surface after vapor CdCl~2~ treatment

at 430°C for the times shown in air ambient with \~10 g/l moisture
content.

For the oxide CdTe~2~O~5~ to form, Te must be preferentially removed
from the CdTe lattice, shifting the chemical equilibrium far from
cadmium vacancy condition, thereby restricting the formation of acceptor
complexes, and reducing p-type conductivity. We have observed the
following correlation between cell performance and GIXRD data of vapor
CdCl~2~ treated films, drawing on measurements back to 1997: high
efficiency PVD and VT cells (Eff \> 11%) exhibit either no surface phase
or a low quantity of CdO; low efficiency, characterized by poor V~OC~
and FF, curvature in forward bias, and re-trace hysteresis, exhibit the
native oxides CdTe~2~O~5~ and CdTeO~3~. In previous studies, similar
results were found for devices processed with wet CdCl~2~ treatment, and
additional Cd-Cl-O phases were detected. The relative sensitivity of wet
treatment to vapor treatment has not been directly evaluated, but
ambient composition is expected to play a critical role in both; in the
case of the wet process, this process sensitivity may depend on the
ability to form a uniform CdCl~2~ coating. These results suggest that it
is important to minimize the humidity during CdCl~2~ vapor treatment.

### 2.2.3 Transparent Contacts

Transparent contacts facilitate bifacial analysis of solar cells under
illumination. Such analysis permits direct probing of the back contact
junction and will be a critical tool for determining properties and
operation of the back contact and their dependence on processing. In
previous work it has been shown that Cu/ITO contacts and ZnTe: Cu films
are viable candidates for this task and can be substituted directly for
C ink or evaporated metal contacts routinely employed [^2]^,,^[^3].

ZnTe films containing different quantities of copper were deposited by
galvanic electrodeposition on TEC-15 conductive and on First Solar
CdTe/CdS solar cell structures. Two types of bath chemistry were
evaluated with respect to the anion, based on sulfate and chloride. The
chemical bath consisted of a pre-mixed aqueous solution containing 0.1M
ZnSO~4~ or ZnCl~2~ and 0.001M TeO~2~. The total volume was 80 mL, with
72mL of ZnSO~4~ or ZnCl~2~ and 8mL of TeO~2~. The electrochemical cell
consisted of a zinc foil anode in electrical contact with the substrate.
The pH of the solution was maintained at \~ 3 by adding drops of dilute
HCl or H~2~SO~4~.

For the initial survey, deposition was carried out in the sulfate bath
varying only the temperature and the deposition time, from 50°C to 70°C
and from 8 to 30 min. For pH = 2 and for depositions carried out with a
stirred solution, reflective metallic films were obtained. For higher pH
and unstirred solutions, orange-brown deposits were obtained. Based on
the visible uniformity of the deposits and on the optical transmission
spectrum, deposition temperature and time of 65°C and 20 min,
respectively, in an unstirred bath at pH = 3 were adopted as the
standard conditions for further experiments.

The second set of depositions was carried out to evaluate the influence
of copper on the properties of the ZnTe films. This was achieved by
varying the copper concentration in both sulfate and chloride baths from
0 to 5 x 10^-4^ M. Optical transmission was measured on samples
deposited on TEC-15 substrates. The highest transmission was obtained
for \[Cu\] ≤ 5 x 10^-5^ M, resulting in a loss of 2-5% transmission from
900 to 1400 nm. Based on this, films were deposited with \[Cu\] = 5 x
10^-5^ M on CdTe/CdS/TEC-15 films, after CdCl~2~ treatment and BM etch.
GIXRD analysis of the sample deposited in from a chloride bath shows
three reflections in addition to those for CdTe (Figure 10.) Indexing
these peaks suggests the presence of the hexagonal form of ZnTe.
Reduction of the integrated peak area of ZnTe and CdTe provides an
estimate of the ZnTe thickness, which is approximately 50nm. No
additional peaks were found for the films deposited in sulfate bath,
suggesting that either the deposit thickness is thinner than the
detection limit (\~5 nm) or is poorly crystallized.

![](media/image12.png){width="5.116666666666666in"
height="2.7465277777777777in"}

[]{#__RefHeading___Toc469989141 .anchor}Figure 10. GIXRD pattern of
ZnTe:Cu/CdTe/CdS/TEC-15

for \[Cu\] = 5 x 10^-5^ M in chloride bath.

### 2.2.4 High throughput VT CdTe characterization and deposition studies

IEC is continuing to investigate the relationship between CdTe
deposition and film properties for cells with CdTe deposited at a high
rate on moving substrates by (VT). The VT deposition process is similar,
but not identical, to that used at First Solar. The influence of He and
O~2~ flow rate and pressure during the CdTe growth were investigated.
Changes in the CdTe surface structure and chemical state were measured
by AFM, GIXRD and contact wetting angle.

No modifications have been made to the original design of the system.
The ceramic gas manifold was replaced when the original one developed a
crack. The range of key deposition parameters and values used for
baseline device evaluation are listed in Table 6.

[]{#__RefHeading___Toc469989100 .anchor}Table 6. Key deposition
parameters used in depositing VT CdTe films

  ---------------------------------- -------------------- ----------------
  Deposition condition               Experimental range   Baseline value

  Substrate T (°C)                   400-550              500

  Source T (°C)                      840-860              840

  Pressure (T)                       20                   20

  He carrier gas flow (sccm)         10-30                20

  O~2~ source 0.5% O~2~/Helium       0-8                  2
  (sccm)                                                  

  CdTe charge (grams)                5.5-6.5              5.5-6.5
  ---------------------------------- -------------------- ----------------

# 

The majority of the CdTe depositions during this contract period were
baseline depositions. These films were used to optimize device
processing as discussed in Section 2.3.2.

However, one study examined the effect on CdTe grain size (as determined
by AFM) of changing the flow rate of the carrier gas (Helium) while
keeping the same ratio of the oxygen source gas (2% O~2~/He) introduced
into the reactor. We also looked at the result of doubling the ratio of
He to oxygen in our baseline film. A summary of results from this study
is given in Table 7. AFM images are in Figures 11-14. As compared to the
baseline run (VT86) increasing the total flow of helium through the
source, results in a substantial decrease in grain size (VT87).
Decreasing the total flow of helium yields slightly larger, more densely
packed grains (VT88). Increasing the ratio of helium to oxygen by two,
results in larger less densely packed grains (VT89). Device results from
these runs are awaiting resolution of problems discussed below.

[]{#__RefHeading___Toc469989101 .anchor}Table 7. Effect of gas flow on
thickness and growth rate of VT CdTe films. All other parameters were as
listed for baseline films in Table 6. VT86 was deposited at baseline
conditions.

+-----------+-----------+-----------+-----------+-----------+-----------+
| Run       | Helium    | 2         | Pressure  | Thickness | Growth    |
|           |           | %O~2/~/He | (Torr)    |           | Rate      |
|           | (sccm)    | (sccm)    |           | (µ)       |           |
|           |           |           |           |           | (µ/min.)  |
+-----------+-----------+-----------+-----------+-----------+-----------+
| VT86      | 20        | 2         | 20        | 5.3       | 1.3       |
+-----------+-----------+-----------+-----------+-----------+-----------+
| VT87      | 30        | 3         | 20        | 3.8       | 1         |
+-----------+-----------+-----------+-----------+-----------+-----------+
| VT88      | 10        | 1         | 20        | 5.4       | 1.4       |
+-----------+-----------+-----------+-----------+-----------+-----------+
| VT89      | 20        | 1         | 20        | 4.8       | 1.2       |
+-----------+-----------+-----------+-----------+-----------+-----------+

+-----------------------------------+-----------------------------------+
| ![](media/image13.                | ![](media/image14.                |
| jpeg){width="3.115972222222222in" | jpeg){width="3.115972222222222in" |
| height="2.5541666666666667in"}    | height="2.5541666666666667in"}    |
+-----------------------------------+-----------------------------------+
| []{#__RefHeading___Toc469989142   | []{#__RefHeading___Toc469989143   |
| .anchor}Figure 11. AFM image of   | .anchor}Figure 12. AFM image of   |
| VT86                              | VT87                              |
|                                   |                                   |
| **(see Table 7)**                 | (see Table 7)                     |
+-----------------------------------+-----------------------------------+
|                                   |                                   |
+-----------------------------------+-----------------------------------+
|                                   |                                   |
+-----------------------------------+-----------------------------------+
| ![](media/image15.                | ![](media/image16.                |
| jpeg){width="3.115972222222222in" | jpeg){width="3.115972222222222in" |
| height="2.5541666666666667in"}    | height="2.5541666666666667in"}    |
+-----------------------------------+-----------------------------------+
| []{#__RefHeading___Toc469989144   | []{#__RefHeading___Toc469989145   |
| .anchor}Figure 13. AFM image of   | .anchor}Figure 14. AFM of VT89    |
| VT88                              |                                   |
|                                   | (see Table 7)                     |
| (see Table 7)                     |                                   |
+-----------------------------------+-----------------------------------+
|                                   |                                   |
+-----------------------------------+-----------------------------------+
|                                   |                                   |
+-----------------------------------+-----------------------------------+

### 2.2.5 Device results and process studies of high throughput VT CdTe

### 

IEC is continuing to investigate the relationship between device
processing and device performance for cells with VT CdTe deposited at a
high rate on moving substrates. Focus has been on establishing and
understanding baseline processing for the post-deposition procedures of
completed CdTe/CdS structures. For the CdTe/CdS post-deposition
treatments, the time and temperature of the CdCl~2~ treatment, use of a
high temperature anneal (HTA) prior to CdCl~2~ treatment, and spacing
between the CdCl~2~ plate and CdTe surface were investigated.

Substrates were TEC-15 SnO~2~ with a Ga~2~O~3~ high resistance buffer
layer. All CdS and CdTe layers for the device results in this section
were deposited under nominally identical conditions. The CdS films were
deposited by CSD, which produces conformal ultra-thin CdS films with low
particulate density and high utilization of Cd species and is discussed
in Section 2.4.1. CdS was deposited by CSD to a thickness of \~90 nm and
then received a CdCl~2~ treatment at 420°C for 30 min. Following
deposition of a \~4 µm CdTe layer at 550°C in a single 6 min pass in
He/O~2~ at 20 T, pieces receive a HTA and/or CdCl~2~ treatment. Contact
processed consists of a BDH etch, then loaded into the evaporator for
the layered back contact structure of 50 nm Te, 30 nm Cu, 50 nm Au.
Presently, we are using Cu/Au contacts evaporated through a mask since
this is more repeatable, quicker, and less likely to introduce shunts
than scribing around a graphite paste contact, as we have previously
used. However, Cu/Au contacts typically give 20-40 mV lower V~oc~ than
graphite. Post-contacting annealing was investigated, and a 30 minute
anneal in Ar at 190°C was found to be optimum. All devices reported in
this section received this anneal.

Run VT 83 varied the time and ambient of the CdCl~2~ treatment of the
completed CdS/CdTe structure at 430°C. Figure 15 shows dependence of
efficiency on the time of the CdCl~2~ treatment ranging from 0 to 60
minutes in either dry or moist air. All pieces received an HTA of 575°C
for one hour prior to the CdCl~2~ except for one piece, which received
no CdCl~2~ either. (A problem during the e-beam deposition of the Te
layer for the contacts resulted in lower efficiency than expected, but
all samples from VT 83 can be compared to each other for relative
purposes). Comparing the two pieces which had no CdCl~2~ treatment, it
is clear that the HTA by itself accomplishes the same enhancement in
performance, to \~8% efficiency, as the optimum CdCl~2~ treatment in
this experiment. This is the best device we have made without a CdCl~2~
treatment. The QE of the device, which received neither HTA nor CdCl~2~,
indicated poor collection from the CdTe. There was no clear difference
in device performance between dry or moist air suggesting the BDH
treatment rendered the CdTe surfaces equally clean in both cases.
However, before BDH there were differences in the type of oxides
detected by GIXRD and in the surface energy determined by the
contact-wetting angle.

![](media/image17.wmf){width="3.4993055555555554in"
height="3.4993055555555554in"}

[]{#__RefHeading___Toc469989146 .anchor}Figure 15. Dependence of
efficiency on the time of the CdCl~2~ treatment

ranging from 0 to 60 minutes in either dry or moist air for VT 83.

It has been shown, by IEC and others, that O~2~ is critical for the
CdCl~2~ treatment to achieve full impact on the device. This leads to
concerns about sufficient circulation and replenishment of the O~2~
ambient during the CdCl~2~ treatment. We have also shown that humidity
and exposure to high temperatures in presence of air can lead to oxide
formation on the CdCl~2~, leading to concern about reusing CdCl~2~
plates. In run VT 85, the spacing between the CdCl~2~ plate and the
substrate was varied during the CdCl~2~ treatment. Since this is a
static system, increasing the distance would increase the availability
of O~2~ at the CdTe surface. We also investigated the effect of the
freshness of the plates, and the HTA in run VT85. Cells made with fresh
plates included those with 0 and 6X spacing where X is our standard
spacing of \~0.6 mm. After being used for treating the cells with no
spacing, the CdCl~2~ plate was re-used for treating with 2X spacing. The
results in Table 8 show that the larger spacing was better,
demonstrating that the delivery of O~2~ and CdCl~2~ are essential to
obtaining good device performance. The HTA resulted in roughly a 2%
decrease in absolute efficiency as well as lower yield. This detrimental
effect of the HTA contradicts many previous results such as Figure 15.
Devices processed with used CdCl~2~ plates had much lower efficiency
and/or severe retrace instability. Following this, fresh CdCl~2~ plates
were used for every treatment. Figure 16 shows the QE from the 4 devices
with the fresh plate. The 33 pc, which had the largest spacing with a
fresh plate and no HTA, had the highest efficiency and is the only
normal QE. The other 3 show reduced CdS thickness due to interdiffusion
and poor collection from the CdTe. This fall-off in collection between
600 and 850 nm is commonly seen in devices from other groups as well. A
flat QE with no decrease in this region is believed to indicate maximum
collection from the CdTe hence optimum processing. From these results,
we concluded that an in-depth evaluation of the HTA treatment was
needed.

[]{#__RefHeading___Toc469989102 .anchor}Table 8. Efficiency and yield of
pieces with different spacing of CdCl~2~ plate and

HTA from run VT85. All cells on piece 22 had unstable retraces.

+---------+------------+---------+---------+---------------+---------+
| piece   | spacing    | Fresh   | HTA     | Efficiency    | Yield   |
|         |            |         |         |               |         |
|         |            | Plate   |         | (%)           |         |
+---------+------------+---------+---------+---------------+---------+
| -12     | 0          | Yes     | Yes     | 7.5           | 3/8     |
+---------+------------+---------+---------+---------------+---------+
| -23     | 0          | Yes     | No      | 9.4           | 6/8     |
+---------+------------+---------+---------+---------------+---------+
| -22     | 2X         | No      | Yes     | unstable      | n/a     |
+---------+------------+---------+---------+---------------+---------+
| -43     | 2X         | No      | No      | 3.0           | 8/8     |
+---------+------------+---------+---------+---------------+---------+
| -42     | 6X         | Yes     | Yes     | 9.4           | 4/8     |
+---------+------------+---------+---------+---------------+---------+
| -33     | 6X         | Yes     | No      | 11.2          | 8/8     |
+---------+------------+---------+---------+---------------+---------+

![](media/image18.wmf){width="3.2881944444444446in"
height="3.2423611111111112in"}

[]{#__RefHeading___Toc469989147 .anchor}Figure 16. QE at 0V for devices
from VT85 with 2 values CdCl~2~

plate-to-substrate distance 0X and 6X as explained above. Devices with
and

without HTA are shown, efficiency values are in Table 14.

The best device from this series was VT85.33-01. The light and dark J-V
curve is shown in Figure 17, having V~oc~=0.749 V, J~SC~=22.9 mA/cm~2~,
FF=65.2%, and Eff.=11.2%. There is some light-dark crossover, which
suggests that the CdS or CdTe are photoconductive. The QE is shown in
Figure 16.

![](media/image19.wmf){width="3.6527777777777777in"
height="3.513888888888889in"}

[]{#__RefHeading___Toc469989148 .anchor}Figure 17. Light and dark J-V
curve on VT85.33-5.

Several VT runs were made after VT85 to sort out the contradictory
dependence on HTA, effects of gas flows during CdTe deposition, and
contacting. Device performance was universally poor, with a high
incidence of shunting and poor retrace stability. Various
characterization methods indicated the CdTe films were dense,
well-oriented, large grain, and largely pin-hole free films. Thus, bulk
CdTe structural properties were not likely to be responsible. Processing
of First Solar material with IEC's CdCl~2~ and contact yielded expected
results of \~9% efficiency. This indicated the problem was with IEC's
CdS or CdTe deposition.

As shown above, QE data suggested significant consumption of the CdS was
occurring in some devices for some process conditions. Thinning of the
CdS and the resulting interdiffusion can lead to poor device
performance. The QE at short wavelengths is a good indication of CdS
thickness. Figure 18 is scatter plot of the V~oc~ vs the QE at 450 nm
for 40 devices processed recently. There is no correlation, indicating
that while there is an unacceptably large variability in the final CdS
thickness, it is not responsible for the unacceptably large variability
in V~oc~.

![](media/image20.wmf){width="3.7222222222222223in"
height="3.4166666666666665in"}

[]{#__RefHeading___Toc469989149 .anchor}Figure 18. Scatter plot of QE at
450 nm vs V~oc~ for 40 CdTe devices.

The CdS and CdTe depositions were nominally the same for all devices

but the CdS and CdTe processing steps were different.

The filament and gas manifold were replaced. They not only were
potential sources of impurities but were also aging. Device results were
still poor. New runs are planned and samples are being sent to NREL for
SIMS. New CdTe source material, processing of PVD CdTe as a control,
different CdS (not CSD), and other comparisons will help identify the
source of this new problem. This new source of irreproducibility,
shunting and poor device results prevented IEC from establishing a high
throughput process with a 12% baseline process as planned for this year.

## 2.4 Improving CdTe/CdS Solar Cell Performance

### 2.4.1 Alternative CdS Deposition

Optimizing the current density in heterojunction solar cells using CdS
window layers requires minimizing the CdS film thickness to reduce
optical absorption losses for photons at λ\< 500 nm. To ensure spatially
uniform device operation, the CdS film must be continuous and dense and
robust to subsequent device processing. A widely employed technique for
fabricating ultra-thin and dense CdS films, from 30 to 100 nm, is
chemical bath deposition (CBD), whereby CdS films are deposited onto
suitable superstrates in a bath containing cadmium and sulfur species
[^4] [^5].

In CBD processing, the heat required to promote the chemical reaction is
delivered from the bath to the sample surface, producing both
heterogeneous CdS nucleation at the surface and homogeneous CdS
formation in the bath. The reaction is enhanced at the hottest region of
the bath. Thus, for baths heated by thermal jacket or water bath,
significant CdS deposition also occurs on the vessel walls, and for
baths heated with an immersion heater, significant deposition occurs on
the heating element. In addition, the bath must be vigorously stirred to
ensure uniform thermal and chemical mixing and to minimize adhesion of
homogeneously nucleated CdS particulates to the growing film surface.
Finally, the disproportionate volumes of the bath and that required
forming the desired CdS film leads to a high proportion of
cadmium-containing waste. Efforts have been undertaken by various groups
to minimize the bath-to-surface volume with the use of cover plates
[^6], but a clear path for combining large area deposition with high
utilization and growth rate for high conversion efficiencies has not
been demonstrated.

The CSD technique presented here[^7] overcomes these limitations, using
the sample surface as the heat source and utilizing the solution surface
tension to minimize the applied liquid volume. The combination of heat
delivery at the surface with low solution volume results in high cadmium
species utilization. The CSD technique, properties of CdS films, and
device results for high efficiency superstrate CdTe/CdS and substrate
CuInGaSe~2~/CdS solar cells made using CSD CdS window layers, have been
presented recently (3^rd^ WCPSEC.)

In CSD, a solution at ambient temperature containing the desired
reactants is applied to a pre-heated surface, as shown in Figure 19.
Film deposition occurs by heterogeneous nucleation at the surface as
heat is transferred into the applied solution. Heterogeneous nucleation
is favored over homogeneous precipitation by thermally enhanced
reactivity at the warmer growth surface, resulting in a high fraction of
cadmium product in the film and, depending on the substrate,
heteroepitaxial film growth. Heat loss from the solution to the ambient
helps maintain favorable conditions for hetergeneous film growth in the
time frame needed for film deposition, \~ 5 minutes.

![](media/image21.wmf){width="2.9965277777777777in" height="1.325in"}

[]{#__RefHeading___Toc469989150 .anchor}Figure 19. Schematic view of CSD
deposition of CdS films.

Stock aqueous solutions of 0.015 M cadmium sulfate, CdSO~4~, 1.5 M
thiourea, CS(NH~2~)~2~, and 14.28 M ammonium hydroxide, NH~4~OH, were
placed in titrating burets for volumetric dispensation. The solutions
were dispensed into a mixing beaker with deionized water to form the
working solution, at a temperature of 25ºC. The volumetric proportions
were 2.2 ml CS(NH~2~)~2~ and 2.8 ml NH~4~OH, with CdSO~4~ and water
adjusted to control the Cd^++^ concentration. The starting pH of the
working solution is \~11.5. Borosilicate glass substrates coated with
ITO were placed under a glass cover on a heated plate at temperatures of
55ºC, 65ºC, 75ºC, 85ºC, and 95°C, and allowed to heat for 10 minutes
prior to application of the solution. A volume of 0.36 ml of working
solution per square inch was applied to the surface of each substrate.
For \[Cd^++^\] = 3 mM, this volume contains sufficient Cd to form CdS
films \~50 nm thick. At the end of the prescribed reaction times,
ranging from 2-6 min, residual solution was transferred from each sample
surface to a waste beaker. The coated substrates were rinsed in flowing
de-ionized water and dried under a forced Ar gas stream. The time to
complete the reaction was determined by visual inspection of the growth
process and varied with temperature for fixed total quantity of solution
is shown in Table 9.

[]{#__RefHeading___Toc469989103 .anchor}Table 9. Time to complete
reaction.

+-----------------------------------+-----------------------------------+
| Substrate Temp                    | Saturation Time                   |
|                                   |                                   |
| (ºC)                              | (min)                             |
+-----------------------------------+-----------------------------------+
| 95                                | 2                                 |
+-----------------------------------+-----------------------------------+
| 85                                | 2                                 |
+-----------------------------------+-----------------------------------+
| 75                                | 3                                 |
+-----------------------------------+-----------------------------------+
| 65                                | 5                                 |
+-----------------------------------+-----------------------------------+
| 55                                | 6                                 |
+-----------------------------------+-----------------------------------+

CdS film thickness was determined by optical absorption and step
profilometry and was correlated to optical transmission for λ\< 500 nm,
allowing \~1 nm precision in comparative analysis. Using the reaction
times shown in Table 9, films were deposited at different temperatures
and Cd species concentrations (Table 10). At a fixed substrate
temperature, thickness increased with total Cd^++^ concentration until
saturation. This limit arises from loss of ammonia species and stalling
of the reaction as the temperature of the applied solution rises. The
remaining Cd can exist as dissolved Cd^++^ or suspended as sulfide or
hydroxide product.

[]{#__RefHeading___Toc469989104 .anchor}Table 10. CdS thickness (nm) for
different substrate temperatures and Cd^++^ concentration.

  ------------- ------------ ----------- ----------- ----------- -----------
                \[Cd^++^\]                                       
                (mM)                                             

  T (ºC)        1.5          2.0         3.0         4.0         5.0

  55            7                                                

  65            16                                   38          

  75            23           31          42          40          43

  85            23                       50          42          

  95            24                                   31          
  ------------- ------------ ----------- ----------- ----------- -----------

The maximum film thickness and deposition rate were 50 nm and 25 nm/min,
respectively, obtained with 3 mM Cd^++^ at 85ºC. The reaction saturation
times (Table 9), which restrict the thickness and growth rate range,
represent a combination of processes: 1) heat transfer and heterogeneous
reaction at the substrate surface, 2) heat transfer to the solution
resulting in ammonia loss, and 3) heat transfer to the solution and
homogeneous reaction.

A steady state model was developed to provide guidance in understanding
heat transfer from the heating source to the solution. For this, the
liquid film was assumed to act as a boundary layer for conductive heat
transport from the glass to the ambient, with constant power being
supplied to the substrate. The temperature-time profile for the solution
surface is shown in Figure 20 for different initial substrate
temperatures. For a substrate temperature of 80°C maintained at the
hot-plate interface and an ambient temperature of 25°C, the estimated
steady state glass surface temperature is 78°C and the liquid surface
temperature is 76°C. The solution surface heats rapidly, reaching 70°C
in \~1 min and thereafter increases slowly to the steady state
temperature, after \~13 min. The time needed to complete the chemical
reaction (Table 9) is always greater than the time needed to raise the
solution surface temperature (Figure 20), indicating that a significant
fraction of the reaction time is spent at conditions where homogeneous
growth and ammonia loss are favorable. This is expected to impact the
utilization of Cd species in the film.

![](media/image22.png){width="3.959722222222222in"
height="2.8555555555555556in"}

[]{#__RefHeading___Toc469989151 .anchor}Figure 20. Solution surface
temperature versus time for

three initial substrate temperatures and constant applied power.

Utilization of Cd species was calculated by mass balance from the
applied solution to thickness of the deposited CdS film (Table 11) and
was verified in selected cases using atomic absorption spectroscopy of
the remaining surface solution. Cd species utilization was \>80% for
cadmium ion concentrations from 1.5 to 3 mM and increased with reaction
temperature. The loss in utilization at high Cd^++^ concentration is
attributed to the loss in ammonia needed to complex the Cd and may be
recoverable by encapsulating the solution surface during reaction.

[]{#__RefHeading___Toc469989105 .anchor}Table 11. Utilization of Cd^++^
(%)

  ------------- ------------ ----------- ----------- ----------- -----------
                \[Cd^++^\]                                       
                (mM)                                             

  T (ºC)        1.5          2.0         3.0         4.0         5.0

  55            28                                               

  65            65                                   56          

  75            90           85          84          58          51

  85            92                       100         60          

  95            94                                   45          
  ------------- ------------ ----------- ----------- ----------- -----------

Additional CdS film thickness is easily obtained by performing multiple
coatings. For depositions carried out at 65ºC and 80ºC, we verified that
the obtained thickness increased linearly with the number of coatings
for up to 4 coatings. We have used the CSD method to deposit CdS films
with ± 1 nm thickness uniformity on glass substrates up to 15 cm x 15 cm
-± 1 nm size, limited only by the ability to heat the substrate
uniformly, demonstrating the feasibility of CSD for large area coverage.
For large areas, different solution dispensation methods such as
spraying could be effectively utilized, since the applied volumes are
small and adhere to the substrate by surface tension, allowing
deposition in face-up or facedown configurations. The technique is
adaptable for utilizing residual heat from prior processing steps in
cell or module manufacturing and is amenable for flexible substrates and
roll-to-roll operation. We have successfully used the CSD technique to
deposit ZnS and CdZnS films and are presently investigating the
deposition of other buffer layers for thin-film heterojunction solar
cell processing.

CdS films deposited on ITO-coated borosilicate substrates were analyzed
for growth habit, crystallinity, and optical transmission. Surface
morphology was characterized using a Digital Instruments Dimension 3100
AFM in tapping mode with etched silicon tips. Figure 21 compares 1 μm x
1 μm AFM scans of CSD CdS films deposited at \~70°C on recrystallized
ITO and ITO coated with In~2~O~3~. The substrates had lateral grain
sizes of 10 nm and 100 nm, respectively. The root-mean square roughness
and grain sizes measured in the AFM scans indicate that the CdS growth
habit is conformal with the underlying substrate. Wide area surveys, 10
μm x 10 μm, showed that CSD CdS films contain less than 10^5^ adherent
particulates per square cm, compared to \~10^8^ for conventional CBD CdS
films. Examination by optical microscopy in transmission mode showed the
fractional area of pinholes is less than 10^-6^, which is comparable to
conventional CBD and less than for thermally evaporated CdS films.

![](media/image23.png){width="2.397222222222222in"
height="1.8555555555555556in"}

![](media/image24.png){width="2.397222222222222in"
height="1.823611111111111in"}

[]{#__RefHeading___Toc469989152 .anchor}Figure 21. Tapping AFM images of
20 nm thick CSD CdS films deposited

at 70°C on recrystallized ITO (top) and recrystallized In~2~O~3~/ITO
(bottom).

The phase composition of CSD films was determined by GIXRD at incident
beam angle, Ω = 0.5°. Figure 22 shows the pattern for a 50 nm thick CSD
CdS film deposited at 70°C onto recrystallized ITO/borosilicate glass.
The pattern indicates a broad peak with three components, consistent
with the wurtzite phase of CdS.

![](media/image25.png){width="3.171527777777778in"
height="2.1131944444444444in"}

[]{#__RefHeading___Toc469989153 .anchor}Figure 22. GIXRD pattern of 50
nm thick CSD film

deposited at 70°C on ITO/borosilicate glass substrate.

The optical transmission spectra of Corning 7059 borosilicate glass, 200
nm ITO/glass, and CSD CdS/ITO/glass, normalized for reflection, are
shown in Figure 23. Transmission curves for three CdS film thicknesses,
6 nm, 20 nm and 45 nm, are shown. The CdS band edge is apparent at \~520
nm (2.4 eV). At lower wavelength, the transmission is limited by
absorption in the CdS film, determined by the absorption curve and the
film thickness. Above the band edge, the transmission through the
structure is limited by the glass and ITO absorptivity, resulting in
T/(1-R) \> 90%.

![](media/image26.png){width="4.86875in" height="3.138888888888889in"}

[]{#__RefHeading___Toc469989154 .anchor}Figure 23. Normalized optical
transmission of Corning 7059 borosilicate glass,

200 nm ITO/glass, and with three thicknesses of CSD CdS.

Through film electrical conductivity was measured on 50 nm thick CdS
films on ITO by measuring the current-voltage characteristic between 2
mm diameter In dots and the ITO substrate. Linear behavior was obtained
corresponding to conductivities of \~10^-4^ S/cm in the dark and
\~10^-3^ S/cm at 10 mW/cm^2^ of white light through the CdS film.
Slightly higher values were obtained in cross-grain measurements of CSD
CdS films deposited on glass.

CdTe/CdS devices were fabricated in a superstrate configuration using
CSD CdS applied to borosilicate glass coated with transparent conductive
oxides (TCO) and high resistance oxide (HR) layers. In the examples
shown here, different combinations of HR/TCO/glass are shown: 1)
In~2~O~3~/ITO/7059, 2) SnO~2~/SnO~2~/7059; and 3)
Ga~2~O~3~/SnO~2~/soda-lime (Pilkington TEC-15). In all cases, the HR
layer thicknesses and sheet resistances were 100 nm and \~1000 Ω/sq,
respectively. The CdS films were deposited by CSD at 70-80°C using the
solution chemistry already described. Nominally 4 μm thick CdTe absorber
layers were deposited by PVD onto stationary substrates at 0.1 μm/min at
340°C or by VT onto moving substrates at 1 μm/min at 550°C translating
at 1 cm/min. The structures were annealed at 600°C in argon for 10
minutes and then treated in CdCl~2~:O~2~:Ar vapor at 400°C for 20 min.
The back contact was fabricated by forming Cu~2~Te followed by
application of a contact. Current-voltage (J-V) parameters of the best
cells obtained are shown in Table 12.

[]{#__RefHeading___Toc469989106 .anchor}Table 12. CdTe/CdS device J-V
results.

+----------+-----------+-----------+-----------+-----------+-----------+
| CdTe     | HR/       | V~oc~     | J~sc~     | FF        | Eff       |
|          | TCO/glass |           |           |           |           |
| Dep      |           | (mV)      | mA/cm^2^  | (%)       | (%)       |
+----------+-----------+-----------+-----------+-----------+-----------+
| PVD      | SnO~2~/Sn | 790       | 25.0      | 68        | 13.4      |
|          | O~2~/7059 |           |           |           |           |
+----------+-----------+-----------+-----------+-----------+-----------+
| PVD      | In~2~O~3~ | 790       | 21.8      | 70        | 12.0      |
|          | /ITO/7059 |           |           |           |           |
+----------+-----------+-----------+-----------+-----------+-----------+
| VTD      | G         | 770       | 21.2      | 64        | 10.4      |
|          | a~2~O~3~/ |           |           |           |           |
|          | SnO~2~/SL |           |           |           |           |
+----------+-----------+-----------+-----------+-----------+-----------+

### 2.4.2 Optimization of CSD CdS

It is well established that while thinner CdS leads to higher J~sc~,
there can be an unacceptable decrease in V~oc~ and FF, due to
degradation in the junction integrity. This can be due to formation of
pinholes in the CdS or to nearly complete consumption of CdS during
subsequent device processing. In both cases, the incomplete CdS layer
results in formation of a high recombination SnO~2~/CdTe junction, which
shunts the CdS/CdTe junction in parallel, and in formation of a lower
band gap CdS~x~Te~1-x~ junction region. Thus, high efficiency devices
require a robust and durable thin CdS film, which maintains a continuous
layer for uniform junction formation. In this contract we investigated
the relationship between device processing, materials properties and
device performance for cells with CdS films deposited by CSD, which
produces conformal ultra-thin CdS films with low particulate density and
high utilization of Cd species [^8]. Focus has been on establishing and
understanding baseline processing procedures for the CdS/TCO
superstrates and post-deposition procedures for completed CdTe/CdS
structures.

IEC has previously demonstrated using PVD CdS films that a CdCl~2~ heat
treatment of CdS prior to CdTe deposition was critical to establish high
efficiency as CdS thickness decreased. This was related to minimizing Te
and S interdiffusion of small grain CdS during the CdTe CdCl~2~
treatment. Thus, the processing variables investigated here for CdS/TCO
were time and temperature of the CdCl~2~ treatment given to the CSD CdS
prior to CdTe deposition. For the CdTe/CdS post-deposition treatments,
the time and temperature of the CdCl~2~ treatment given to the CdS/CdTe,
and use of a high temperature anneal (HTA) prior to CdCl~2~ treatment
for both CdS and CdTe were investigated, and discussed further in
Section 2.3. Substrates were TEC-15 SnO~2~ with a Ga~2~O~3~ high
resistance buffer layer. All CdS and CdTe layers were deposited under
nominally identical conditions for all devices presented in this
section. CdS was deposited by CSD to a thickness of \~90 nm. The CdS was
given different treatments prior to CdTe deposition. Following
deposition of a \~4 µm CdTe layer at 550°C in a single 6 min pass in
He/O~2~ at 20 T, pieces may receive a high temperature anneal (HTA)
and/or CdCl~2~ treatment. Contact processing consists of a BDH etch then
loading into the evaporator for 50 nm Te, 30 nm Cu, 50 nm Au. Presently,
we are using Cu/Au contacts evaporated through a mask since this is more
repeatable than scribing around a graphite paste contact but typically
leads to 20-30 mV lower V~oc~. Samples receive a 30 minute anneal in Ar
at 190°C before testing.

Figure 24 shows optical transmission data the TEC15 superstrate, TEC15 +
40 nm Ga~2~O~3~ high resistance layer, and a 90 nm thick CdS film before
and after vapor CdCl~2~ treatment at 430°C. Evidence of significant
recrystallization is indicated by AFM profiles of the surface, (Figure
25) and sharpening of diffraction peaks on GIXRD patterns.

The influence of the Ga~2~O~3~ high resistance layer on CSD growth was
studied. It was found that CdS films grown on Ga~2~O~3~ or In~2~O~3~ HR
layers exhibit larger grain size than those on glass or ITO, due to the
conformal growth habit and larger grains obtained in the HR films.

![](media/image27.png){width="4.899305555555555in"
height="3.1597222222222223in"}

[]{#__RefHeading___Toc469989155 .anchor}Figure 24. Normalized optical
transmission, T/(1-R), of window layer used in VT cells.

+-----------------------------------------------------------------------+
| ![](media/image28.png){width="2.7020833333333334in"                   |
| height="2.591666666666667in"}                                         |
+-----------------------------------------------------------------------+
| ![](media/image29.png){width="2.801388888888889in"                    |
| height="2.591666666666667in"}                                         |
+-----------------------------------------------------------------------+
| []{#__RefHeading___Toc469989156 .anchor}Figure 25. AFM scans of CSD   |
| CdS film: as-deposited (left)                                         |
|                                                                       |
| and after vapor CdCl~2~ HT at 430°C (right).                          |
+-----------------------------------------------------------------------+

Having established how vapor CdCl~2~ treatments change the CdS
properties, we attempted to optimize the device performance by varying
the temperature and CdCl~2~ treatment time. Run VT 79 exposed identical
CdS layers to CdCl~2~ for various temperatures (350-430°C) and times
(5-20 min). After CdTe deposition, all pieces but one received a 30 min
HTA at 580°C in Ar and all received a 25 min CdCl~2~ treatment at 400°C.
Figure 26 shows the efficiency of the best cell per piece for various
times and temperatures of the CdS treatment. In general, efficiency
improved with increasing the temperature and time. Therefore, a
temperature of 430°C and time of 20 min was selected for CdS treatment
in subsequent studies. The lone piece, which did not receive the
post-CdTe HTA, had noticeably poorer efficiency.

![](media/image30.wmf){width="3.08125in" height="3.196527777777778in"}

[]{#__RefHeading___Toc469989157 .anchor}Figure 26. Effect of CdCl~2~
treatment time of CdS layers on

efficiency at different temperatures from run VT 79.

Since the CdS is deposited from a Cl based-solution, it was possible
that an anneal alone would be sufficient to activate the CdS grain
growth against further interdiffusion, thus, eliminating the need for a
separate CdCl~2~ treatment of the CdS. It was decided to evaluate the
application of an HTA step to only the CSD CdS layer.

Run VT 84 compared devices with and without the HTA and CdCl~2~
treatments to CdS prior to CdTe deposition. The CdS CdCl~2~ treatment
was at 430°C for 20 min and the HTA was at 550° for 30 min in Ar. Also
included in this run were two substrates of glass/SnO~2~/CdS from First
Solar, which did not receive any treatments since the CdS was deposited
at high temperature. After CdTe deposition, all IEC pieces received a 20
min HTA at 560°C in Ar followed by a 25 min CdCl~2~ treatment at 420°C
in air. Table 13 lists J-V parameters and yield for this set of devices.
The HTA alone or in combination with CdCl~2~ yielded lower FF. The best
cell on three identically processed pieces with only CdCl~2~ ranged from
10-11%. Yields were better than 75%.

Devices on First Solar CdS had similar V~oc~ and FF to IEC's VT cells.
This was investigated further by analyzing the dark J-V curves from the
best device on 33 (IEC CdS) and 43 (FS CdS) pieces, which had similar
V~oc~ and FF. Both devices had diode factors A\~2, recombination current
density J~o~\~1E-5 mA/cm^2^, and R~S~\~ 2 -cm^2^. We conclude that
IEC's CdS is not a major limitation to higher V~oc~ or yield at this
time providing it has a CdCl~2~ pre-treatment.

Figure 27 shows the QE curves measured at --1V with light bias on
devices from 4 pieces of VT 84 from Table 13. Below 550 nm, they fall
into three groups. The FS CdS is thicker so it has a much lower
transmission and QE. Devices with IEC CdS which received either HTA *or*
CdCl~2~ have QE values expected for \~90 nm of CdS. But the sample
receiving both HTA *and* CdCl~2~ treatments shows evidence of S-Te
interdiffusion[^9]; i.e. the QE has increased at both short and long
wavelengths compared to the other pieces. Note the rounding of the QE
between 500 and 600 nm for the --13 piece (HTA without CdCl~2~). This
also suggests formation of CdS~x~Te~1-x~. Based on these results we have
selected CdCl~2~ at 430°C for 20 min without HTA as the standard process
for device fabrication with CSD films.

In summary, device performance exhibited sensitivity to the CSD CdS
treatment, similar to PVD CdS cells. The vapor CdCl~2~ ambient
composition during treatment of CdTe/CdS structures was shown to affect
cell performance, and is attributed to the vapor concentrations of
CdCl~2~, O~2~ and H~2~O. The CdCl~2~ treatment sharpens the optical
transmission edge and improves crystallinity of CSD CdS films just as
with PVD CdS films.

[]{#__RefHeading___Toc469989107 .anchor}Table 13. Performance of best
cell-on-pieces from VT 84 having differently treated CdS. Samples CdS
from FS included. Yield based on shunting and low FF of 8 cells per
piece.

+--------+---------------+----------+----------+-------+-------+-------+
| Piece  | CdS treatment | V~oc~    | J~sc~    | FF    | Eff.  | Yield |
|        |               |          |          |       |       |       |
|        |               | \(V\)    | (m       | (%)   | (%)   |       |
|        |               |          | A/cm^2^) |       |       |       |
+--------+---------------+----------+----------+-------+-------+-------+
| -12    | HTA + CdCl~2~ | 0.640    | 19.8     | 41.2  | 5.2   | 6/8   |
+--------+---------------+----------+----------+-------+-------+-------+
| -13    | HTA only      | 0.740    | 21.2     | 58.4  | 9.1   | 6/8   |
+--------+---------------+----------+----------+-------+-------+-------+
| -22    | HTA only      | 0.718    | 22.0     | 60.4  | 9.5   | 8/8   |
+--------+---------------+----------+----------+-------+-------+-------+
| -23    | CdCl~2~ only  | 0.773    | 22.3     | 63.2  | 10.9  | 7/8   |
+--------+---------------+----------+----------+-------+-------+-------+
| -32    | CdCl~2~ only  | 0.742    | 21.9     | 61.9  | 10.0  | 8/8   |
+--------+---------------+----------+----------+-------+-------+-------+
| -33    | CdCl~2~ only  | 0.747    | 22.3     | 65.6  | 11.0  | 7/8   |
+--------+---------------+----------+----------+-------+-------+-------+
| -42    | FS CdS        | 0.720    | 20.5     | 63.5  | 9.4   | 7/8   |
+--------+---------------+----------+----------+-------+-------+-------+
| -43    | FS CdS        | 0.743    | 20.4     | 67.0  | 10.2  | 8/8   |
+--------+---------------+----------+----------+-------+-------+-------+

![](media/image31.wmf){width="3.2881944444444446in"
height="3.234722222222222in"}

[]{#__RefHeading___Toc469989158 .anchor}Figure 27. QE of 4 pieces from
VT 84 as described in Table 13.

The 43 piece had FS CdS while the other three had IEC CdS.

### 2.4.3 Microscale Junction Analysis

We are investigating the use of electric, or electrostatic, force
microscopy as a tool for directly measuring field profiles and carrier
transport in operational thin-film CdTe/CdS solar cells. Different
measurement modes are possible depending on the electrical
configuration; we have elected to investigate cells in cross-section
using a conductive tip to measure potential or current under different
light and voltage bias conditions. Figure 28 below shows the physical
and electrical configuration of the solar cell and microscope for
electric field gradient imaging using interleaving. During operation,
the measured voltage is obtained with respect to the morphology
displacement voltage by interleaving the topological and potential
measurements.

![](media/image32.wmf){width="6.470138888888889in"
height="3.782638888888889in"}

[]{#__RefHeading___Toc469989159 .anchor}Figure 28. Schematic diagram
(not to scale) of taper mounted

solar cell and microscope arrangement for potential measurement.

Electric force mode, with bias from solar cell to tip is also being
used. The spatial resolution of this technique, combined with electronic
models of device operation, may enable micro-scale variations to be
correlated with processing variations and macro-scale device
performance. For this, a method for reliably mounting and grinding and
polishing electrically active CdTe/CdS cells in cross-section was
developed, and a Digital Instruments Nanoscope AFM has been modified for
voltage and current measurement. The completed cell structure has
external electrical leads connected to each side of the junction and is
mounted so that illumination bias can be provided. Figure 29 shows the
sample mount configuration. Figure 30 shows a 2 m x 2 m tapping mode
image of a First Solar cell polished at a 45 degree angle.

![](media/image33.wmf){width="3.0305555555555554in"
height="1.6256944444444446in"}

[]{#__RefHeading___Toc469989160 .anchor}Figure 29. Schematic of solar
cell mounted in epoxy at an angle, exposing the junction.

![](media/image34.wmf){width="4.147222222222222in"
height="3.9923611111111112in"}

[]{#__RefHeading___Toc469989161 .anchor}Figure 30. Tapping AFM image of
the CdTe-CdS interface region

of a thin-film First Solar CdTe/CdS solar cell after mechanical
polishing.

Based on J-V measurements of the cell after mounting and polishing, the
fracturing, mounting and polishing process does not alter the
current-voltage behavior of the cell. Furthermore, preliminary
measurements made in voltage mode in the dark at forward bias show
detectable contrast differences between CdS and CdTe sides of the
junction. However, interpretation of the contrast in these measurements
will ultimately depend on relative changes measured in systematic
studies. This technique will be evaluated on cells fabricated with
different post-deposition processing and before and after stress-induced
changes.

### 2.4.4 CdTe Conductivity and Alternative Treatments

As-deposited CdTe films do not have the required electronic properties
needed to obtain high quality junctions. Unlike CuInSe~2~ and related
materials, the extremely narrow single-phase region coupled with the
high constituent element vapor pressures leads to stoichiometric
deposits and low as-deposited film conductivity. Thus, post deposition
treatments are needed to shift the defect equilibrium sufficiently to
allow formation of acceptor complexes, which have relatively low
ionization potentials. As Marfaing has shown for single crystal CdTe, a
principal result of the CdCl~2~ treatment in oxygen containing ambient
is production of V~Cd~ via removal of lattice-bound Cd and formation of
CdO [^10]. Analogous behavior may be expected in thin films, and indeed,
at elevated (20 vol %) oxygen levels, CdO product is found on the
surface of CdTe films treated in CdCl~2~ vapor. A proposed
phenomenological model for controlling CdTe conductivity via post
deposition treatments is: 1) formation of V~Cd~ at the surface; 2)
redistribution of the vacancies in the course of the treatment; and 3)
addition of extrinsic dopant to form the desired acceptor complex. This
model offers an explanation of why thin film CdTe cells made in
superstrate mode give superior junction performance than those substrate
mode. Experiments were conducted to explore this model, to improve our
understanding of the treatments, and to find paths to more robust and
controllable alternative treatments.

The initial experiment was to measure the dark and light cross-grain
conductivity versus T on CdTe films from a single PVD deposition
(#41184). The first round of treatments were: none; HTA ≡ Ar \@550°C for
60 min; and Te vapor \@550°C for 60 min. No CdCl~2~ or Cu exposure was
applied to this first group. Several pieces were held at each of these
treatments and subsequent samples from each set will receive CdCl~2~ and
contact treatments as well. Two contacts were used. One piece each
received a 50 nm Au contact and one piece received 50 nm Te/30 nm Cu/50
nm Au contacts. IEC's co-planar contact mask was used, having 0.05cm
spacing. Substrates were all Corning 7059 glass. Conductivity was
measured as a function of temperature in the dark and then at 300K in
the light. Temperature was 220 to 320 K on one set and 210K to 360K on
the second set. N~2~ purging was used. The voltage range was selected to
give measurable currents, typically ±10, 20, 50 V.

Results are in Table 14, showing dark and light conductivity at 300K and
dark conductivity activation energy. Most samples had a curved Arrhenius
plot but could be approximately fit with a straight line. All samples
had linear J-V relation (Ohmic) in the dark and most also in the light.
A few had moderately non-Ohmic behavior in the light but best fit was
used (regression coefficient of at least 0.96). Hot probe measurements
were also made, using a hot tip at 250°C and a cold tip at room
temperature cm as-deposited samples showed slight negative voltage --1
to --10 m V, which is consistent with intrinsic conductivity. After the
anneal in Ar or Te vapor, the samples showed moderate positive voltage
(+1 to +100 mV) consistent with p-type conductivity.

[]{#__RefHeading___Toc469989108 .anchor}Table 14. Conductivity results
for treated CdTe films.

+--------+----------------+---------+-------+--------+-------+-------+
| Piece  | Treatment      | Contact | Dark  | Light  | Ratio | Ea    |
|        |                |         | cond. | cond.  |       |       |
|        |                |         |       |        | L/D   | (eV)  |
|        |                |         | (     | (S/cm) |       |       |
|        |                |         | S/cm) |        |       |       |
+--------+----------------+---------+-------+--------+-------+-------+
| 32a    | None           | Au      | 4E-9  | 8E-7   | 200   | 0.64  |
+--------+----------------+---------+-------+--------+-------+-------+
| 11a    | Ar,550°C,60min | Au      | 7E-7  | 2E-4   | 280   | n/a   |
+--------+----------------+---------+-------+--------+-------+-------+
| 21d    | Te,550°C,60m   | Au      | 6E-7  | 9E-6   | 15    | 0.22  |
+--------+----------------+---------+-------+--------+-------+-------+
| 21c    | None           | T       | 9E-9  | 1E-6   | 110   | 0.37  |
|        |                | e/Cu/Au |       |        |       |       |
+--------+----------------+---------+-------+--------+-------+-------+
| 32b    | Ar,550°C,60min | T       | 1E-7  | 9E-5   | 900   | 0.49  |
|        |                | e/Cu/Au |       |        |       |       |
+--------+----------------+---------+-------+--------+-------+-------+
| 11b    | Te,550°C,60m   | T       | 7E-7  | 5E-6   | 7     | 0.21  |
|        |                | e/Cu/Au |       |        |       |       |
+--------+----------------+---------+-------+--------+-------+-------+

Two clear conclusions emerge from this experiment. First, two untreated
samples (32a and 21c) have the lowest dark conductivity independent of
contact. Second, both samples treated in Te have the lowest
photoconductive ratio, and lowest activation energy. These are
consistent: the higher the Fermi level, the higher the density of gap
states it resides in, so the less its position can be influenced by
trapped charge. Thus, this data is consistent with the premise that Te
makes CdTe more conductive. There is no clear difference between samples
with Au or Te/Cu/Au contacts. In fact, 2 samples with Au contacts had
more linear light J-V than those with Te/Cu/Au, which contradicts the
assumption that Te/Cu is needed to give an Ohmic contact. The next step
in this experiment is vapor CdCl~2~ treatment of the annealed films with
varying concentration of CdCl~2~, oxygen and water vapor.

Alternative approaches for delivering CdCl~2~ to the CdTe surface and
evaluation of the process sensitivity are being investigated in part
under the team activity and in-part as a critical aspect of improving
the post-deposition process to improve yield and performance. For vapor
CdCl~2~ treatment, mathematical models of the delivery of molecular
CdCl~2~ species have been generated for diffusion from a packed bed of
CdCl~2~, for transport via carrier gas *across* a packed bed of CdCl~2~
and for transport via carrier gas *through* a packed bed of CdCl~2~
pellets. For the last case, momentum transfer from the carrier gas to
the pellets provides an additional term to the residence-time/diffusion
term, allowing relatively high CdCl~2~ concentrations to be achieved at
temperatures below 300°C.

Alternative treatment chemistries are also being examined to exploit the
proposed model that the role of the treatment is to shift the electronic
level stoichiometry of the CdTe film. ZnCl~2~ and TeCl~4~ are each
reactive halide compounds with relatively high vapor pressures and which
sublime congruently at lower temperatures than CdCl~2~. The vapor
treatment reactor described elsewhere [^11] is being used to expose CdTe
films to controlled concentrations of these vapors at different reaction
temperatures. For PVD films, grain growth has been observed using both
halide species. For First Solar VT films, devices with efficiency
approaching 9% were obtained using ZnCl~2~ vapor.

### 2.4.5 Effects of Cu as an impurity in CdS

There is little direct control over the CdS electrical properties during
deposition or subsequent processing in CdS/CdTe solar cells. CdS
photoconductivity (PC) was investigated extensively in the 1970s with
regard to Cu-related acceptor states in CdS of CdS/Cu~2~S solar cells
[^12] [^13]. Cu creates deep hole traps in the CdS, greatly increasing
the dark resistivity, but this could be recovered with illumination. It
has been shown that Cu is present in CdS of CdS/CdTe solar cells, even
in special devices made without any intentional Cu layer or Cu doping
[^14] [^15]. The Cu concentration in the CdS also increases after
accelerated thermal and voltage bias stressing Error: Reference source
not found. Several recent papers have investigated the PC in unstressed
CdS/CdTe devices using AQE at far forward bias \~1.5V [^16] [^17] [^18].
We used the bias light spectrum to characterize the AQE in CdS/CdTe
devices from several groups.

The CdS/CdTe devices came from BP Solar, Colorado State University
(CSU), First Solar, University of Toledo (UT), and the IEC. They all had
initial efficiencies of 10-12%. The devices all had some intentional Cu
introduced during processing of the back contact. Two unique devices
made at IEC, one without a CdS layer and one with a Cu-doped CdS layer,
were included. AQE measurements were made under various filtered light
bias conditions. 'Blue' bias light was obtained with a Corning 4-71
filter passing wavelengths below 600 nm and 'red' bias light was
obtained with a Corning 2-58 filter passing wavelengths greater than 650
nm. The probe beam for AQE measurements was chopped at \~75 Hz unless
specified. J-V measurements were also made using filtered light.

Figure 31 shows the AQE for an IEC device at +0.6 V bias with white,
red, blue and no (dark) light bias. Red bias light increases the blue
AQE. This enhancement was evident even at 0 V bias but smaller in
magnitude. With red bias, the AQE exceeds unity below 550 nm. With blue
or white bias, the AQE increases beyond 600 nm compared to the dark.
Qualitatively, all CdS/CdTe devices we have examined (First Solar, BP
Solar, IEC, CSU, UT) at forward bias using color bias light are
relatively similar to results in this figure.

The effect of red bias light can be seen more clearly when plotted as
the ratio of AQE with red bias to AQE with white bias light. Figure 32
shows this ratio for unstressed devices from the five groups. This shows
the effect of eliminating blue light from the bias spectrum. To first
order, all devices show qualitatively similar behavior with an
enhancement in the AQE for λ\<550 nm with red light. The exception is
for a device made at IEC without a CdS layer (measured at 0.1 V since
V~OC~ was only 0.2 V), which showed no enhancementMany different
methods of depositing the CdS, the CdTe and the contact are represented,
yet they all show relatively large enhancement of the blue AQE with red
bias light.

![](media/image35.wmf){width="2.5701388888888888in"
height="2.5083333333333333in"}

[]{#__RefHeading___Toc469989162 .anchor}Figure 31. AQE of IEC CdS/CdTe
device at +0.6 V bias with 4 bias light conditions

![](media/image36.wmf){width="3.3in" height="3.2006944444444443in"}

[]{#__RefHeading___Toc469989163 .anchor}Figure 32. Ratio of AQE with red
bias to AQE with white bias for various CdS/CdTe devices plus one CdTe
device without CdS. Data from First Solar divided by 10.

![](media/image37.wmf){width="3.359722222222222in"
height="3.2319444444444443in"}

[]{#__RefHeading___Toc469989164 .anchor}Figure 33. Ratio of AQE near
V~MP~ with red light to white light for

a standard device, one with Cu-doped CdS and one after it was stressed.

IEC fabricated a device with a 0.5 nm Cu layer on the TCO before
depositing the CdS. Device performance was poor: V~OC~=0.47 V, J~SC~=15
mA/cm^2^, FF=30% with rollover in forward bias and leakage in reverse
bias. These symptoms are also seen to a lesser degree in devices after
extended stressing. (More description of the device and its performance
was given in IEC's annual report for 2002.) Figure 33 shows the ratio of
AQE with red to white light bias for the device with Cu-doped CdS and
one with standard CdS, both measured near their maximum power point.
Deliberate Cu doping leads to greater enhancement of the blue AQE.

The AQE ratio with red to white bias of a stressed device (@V~OC~, 30
days, 65°C, light, air) is also shown in Figure 33. The efficiency
decreased from 10 to 8% after stress. The increase in AQE below 550 nm
with red bias becomes even more pronounced after stress. We have seen
this increase in AQE below 550 nm after stress in devices from several
groups.

The effect of the chopping frequency is shown in Figure 34. The
enhancement decreases with increasing frequency. At 753 Hz, the
enhancement has decreased significantly; the AQE with red bias is very
similar to the AQE with white bias at 73 Hz. A frequency response
cut-off of \~1000 Hz at 28°C indicates that the energy of the deep state
responsible for the capture or emission is \~0.55-0.60 eV from the
respective band edge.

The ratio of AQE with blue bias light to no bias is shown in Figure 35
for some of the same devices as in Figure 32. There is an increase in
AQE with increasing wavelength beyond 550 nm. It also increased by a
factor of \~2-4 with stress and with Cu-doping of the CdS. However, the
AQE with blue bias light was frequency independent.

![](media/image38.wmf){width="3.329861111111111in"
height="3.201388888888889in"}

[]{#__RefHeading___Toc469989165 .anchor}Figure 34. AQE of IEC device
with red bias at various chopping frequency.

AQE with white bias light shown for comparison.

![](media/image39.wmf){width="3.2888888888888888in"
height="3.204861111111111in"}

[]{#__RefHeading___Toc469989166 .anchor}Figure 35. Ratio of AQE with
blue bias to no bias for same devices as Figure 32.

These results can be self-consistently explained as follows. With blue
light (above the CdS band gap), holes are preferentially trapped leading
to an increase in positive space charge. This is the source of the CdS
photoconductivity. Consequently, the depletion width in the CdTe must
increase to maintain space charge neutrality, thus enhancing the
collection of carriers in the CdTe generated by red light. Enhancement
in collection from more deeply penetrating photons is consistent with
the shape of Figure 35. Thus, due to photoconductive nature of CdS, any
bias light containing blue photons will enable increased collection from
red photons absorbed in the CdTe.

This same phenomenon is responsible for the increase in blue AQE with
red bias light except now it is the blue ac chopped light which is
modifying the CdS conductivity and CdTe space charge width. When the
chopped light is off, the CdS conductivity is low with dark or red bias
light, since there is no blue light present. When the blue chopped light
is on, the CdS photoconductivity and CdTe depletion width increase as
described above. In the specific case of red bias light, this enables
enhanced collection of bias-light generated carriers from the CdTe in
synchronization with the blue chopped light. Thus, carriers from the red
bias exit the device in phase with those from the blue chopped light
resulting in the appearance of collecting more than one carrier per
incident blue photon. This is a form of photogating [^19] ^,^[^20].
There is no enhancement with dark bias because there are no
bias-generated carriers to be collected from the CdTe. The AQE\> with
red bias was accompanied by a large increase in magnitude of the phase
angle between real and imaginary components of the ac signal for λ\<550
nm. Others have reported this as an indication of photoconductive gain
in CdS Error: Reference source not
found^,Error:\ Reference\ source\ not\ found,Error:\ Reference\ source\ not\ found^.
There was no change in phase between white, blue or dark bias
conditions.

The effect of series resistance on AQE including the impact of
photoconductive layers has been used to study voltage bias dependent AQE
[^21]. Based on the standard circuit model of a solar cell, a constant
series resistance will cause a wavelength independent *decrease* in the
AQE while a photoconductive resistance will give a wavelength dependent
*increase* in the AQE. Small changes in the operating point of the
device on the J-V curve result from the series resistance, which changes
the current flow. The model assumes that there is no voltage dependent
collection, which is valid for CdS/CIS or c-Si but not CdS/CdTe devices.
The analysis of reference Error: Reference source not found is not
applicable to interpret these bias light dependent measurements. We have
verified that the forward biased QE of CdTe with white light is
determined by the voltage dependence of the photocurrent and therefore,
not likely to be determined by this resistance effect.

In summary, these results are consistent with Cu-doped photoconductive
CdS whose carrier density is modulated by above-band gap (blue) light
from either the ac probe or dc bias light. This is an ac effect and has
little relevance to the performance of a solar cell. The CdS resistance
makes an insignificant contribution to the overall series resistance of
the device under white light, which restores the CdS to a highly
conductive state. PC-CdS is responsible for light-dark crossover.

## 2.5 National CdTe Team Activity and Collaborations

Steve Hegedus and Brian McCandless attended the National Team Meeting
held at NREL on October 31 and November 1, 2002. Brian presented a
summary of stress degradation effects on CdTe cells using data contained
in Team Minutes of meetings to 1996. Brian also presented data on
effects of post processing and anneal and CdCl~2~ treatment on BP Solar
films. Steve presented talks on photoconductivity in CdS films and also
on the effects of ambient on CdTe cells during stress treatments.

Kevin Dobson and Brian McCandless attended the National Team meeting
held at NREL on July 10 and 11, 2003. Two presentations were given,
detailing humidity effects during post-deposition processing and stress.
IEC is represented on sub-teams, by Kevin Dobson, Steve Hegedus and
Brian McCandless. Brian McCandless visited First Solar September 23,
2003. IEC has sustained regular interaction with First Solar in the
analysis and development of the back contact process. This has included
characterization of surfaces by GIXRD, cross-section analysis by AFM,
surface energy measurements, and compositional analysis of alternative
contact materials.

#  

# 

# **3 CuInSe~2~-BASED SOLAR CELLS**

## 3.1 In-line deposition Introduction

## 

Research during this period was primarily concentrated on trying to
resolve the issue of Cu(InGa)Se~2~ film cracking and its adhesion to Mo
in roll-to-roll deposition on Mo coated polyimide (Upilex) web.
Secondarily, a number of changes have been made to the reactor and to
its operation in order to improve controllability and robustness.

To a certain degree the adhesion problem has been solved by depositing a
100 to 300 Å -(Ga~1-x~In~x~)~2~Se~3~ . However, film cracking,
illustrated in the micrograph of Figure 36, has remained untractable.
Figure 36 seems to indicate that cracks on the Cu(InGa)Se~2~ film result
from the cracks that are formed on the Mo film. These cracks on Mo
develop in the Cu(InGa)Se~2~ chamber prior the deposition of the
absorber as can be surmised by the texture difference between the film
just around the cracks and the film away from the cracks. If the
bi-layer Mo/ Cu(InGa)Se~2~ cracked after the deposition such texture
difference would not have been observable. It should be emphasized that
cracking of the Mo films have never been observed prior to the
Cu(InGa)Se~2~ deposition.

![](media/image40.jpeg){width="3.6333333333333333in"
height="2.7069444444444444in"}

[]{#__RefHeading___Toc469989167 .anchor}Figure 36. Cracking of Mo/CIGS
structure on Upilex substrate;

Note the cracks on the Mo film imaged on the Cu(InGa)Se~2~ film.

Based on these observations, the main approach used to investigate film
cracking was to deposit Cu(InGa)Se~2~ films on Mo deposited under
different conditions and on different substrates at the facilities of
Techni-Met, Inc., in Windsor, Connecticut. The cracking was evaluated by
visual observation and the adhesion by the 40 oz/inch tape test. The
parameters that were varied during Mo depositions were substrate
temperature, sputtering pressure both influencing film stress. Two types
of Upilex substrate were evaluated: Upilex S and Upilex SG, the latter
providing better adhesion to metal films, according the manufacturer.
These substrates were used as is or after a surface conditioning oxygen
plasma treatment. The web was either cooled by a backing drum chilled to
12°C or its temperature was not actively controlled and reached closed
to 100°C from the effect of the sputtering plasma. In addition to
testing Upilex substrates coated with Mo deposited at different
conditions, the effect of the web tension during Cu(InGa)Se~2~
deposition was also evaluated.

### 3.1.1 Cracking and Adhesion

A number of Upilex/Mo substrates were evaluated in terms of cracking
propensity during CuInGaSe~2~ deposition and Mo/CuInGaSe~2~ adhesion.
All the samples evaluated showed some degree of cracking. The
CuInGaSe~2~ adhesion was found to be variable, being acceptable in some
cases i.e.; less than 10% film removal with the 40 oz/inch tape, and
heavy flaking in the others.

There was no correlation between any Mo sputtering parameters and the
level of cracking or adhesion. This was the case for both Upilex S and
Upilex SG. Oxygen plasma treatment of the Upilex surface prior to Mo
deposition didn\'t result in any measurable effect. The lack of
correlation was also observed between cracking and adhesion; and the web
tension during the CuInGaSe~2~ deposition for tension levels between 8
lb to slightly less than 2 lb for 6\" wide web used. Investigation in
this area is continuing with the focus on both the Mo film and the
process of deposition of the CuInGaSe~2~ film.

### 3.1.2 System Improvements

The system\'s Cu and In sources had been initially operating by Atomic
Absorption Spectroscopy (AAS) control while thermocouples (TC) were used
only for monitoring purposes. It was decided to operate these sources by
TC control and use AAS as the monitoring tool in order to better
understand the systems operational characteristics and also judge the
reproducibility under TC control. Under AAS control reproducibility was
marginally acceptable as the alignment of the units\' optical components
were found to require frequent adjustments. The initial TC control
experiments showed large fluctuations in the power delivered to the Cu
source, which operates at a temperature of around 1380°C, considerably
higher than the other sources. These power fluctuations in effect
rendered the system uncontrollable. Considerable amount of time was
spent to trace the cause of the problem that was traced to the noise in
the TC extension wires that National Instrument Field Point system was
not able to filter. As a result all the source TC extension wires were
replaced by shielded wiring and the shields were connected to the
floating ground of the Field Point modules. Tests with the new wiring
system didn\'t display any power fluctuations resulting in proper source
control.

Following the shielding of TC extension wiring, the system went through
some testing since it has also undergone a number of changes and parts
replacements. As a result of these tests it became apparent that the TC
control of the sources with AAS monitoring was found to give better
reproducibility and robustness. However, since a switch was made from
naked TCs to Ta sheathed TCs an offset existed between the readout and
actual temperature. In the case of Cu source, this offset was determined
by observing the temperature plateau during the cool down, which
corresponds to the solidification point of Cu i.e. 1085°C. For the Ga
and In sources the offsets were determined experimentally by measuring
the film composition and estimating from it the adjustment to the source
temperatures as obtained from the effusion model of the system.

Another problem observed during the operation of the system was the
controllability of the Se sparger in that at operational temperatures Se
sparger heater power supply was found to draw only 1% power. This
indicated that Se source was primarily heated by radiation from other
heat sources in the system. In order to remedy this situation various
heat shielding schemes have been tried. As a result, Se source
temperature was reduced to 320°C from 360°C without any change in the
film composition and structure.

There was, however, a need to make Se effusion a totally independent and
controllable parameter. The Se sparger system consists of a source
bottle heated by an electrical heater, a main manifold and four branches
with two effusion holes each (Figure 37) along the sides of the Cu, Ga
and In sources.

![](media/image41.jpeg){width="5.069444444444445in"
height="2.4305555555555554in"}

[]{#__RefHeading___Toc469989168 .anchor}Figure 37. Source layout in the
CIGS in-line deposition system.

The primary issue was the control of the surface temperature of the Se
melt in the source bottle. Since Se source receives substantial amount
of heat from Cu, Ga and In sources a heat shield was installed between
these sources and the Se bottle. There is however heat coming into the
Se bottle by conduction from the manifold. This heat source along with
the spiral electric heater placed towards the top of the Se bottle;
create a temperature gradient in the Se source. Se temperature is the
highest at the top where the manifold is attached, and the coldest at
the bottom of the source. Under these circumstances as Se is used up the
level goes down and since the control thermocouple immersed into the
melt is stationary, it indicates the true surface temperature only when
the melt level is at the same level as the thermocouple tip. This
problem was solved my using a 0.01\" OD inconel sheeted thermocouple
attached to a graphite ring placed inside the source with a certain
amount of slack on the thermocouple. During the operation the graphite
ring floats over the melt, and since the thermocouple is thin, as the
melt level changes graphite ring stays floating over the melt and the
thermocouple always give the surface temperature as long as there is
source material in the bottle.

Another problem with controlling Se flux onto the substrate was that Se
was found to condense in the branches during start-up and/or cool-down
creating secondary uncontrolled sources. This problem was solved by
carefully monitoring the temperature of all the source shields, Se
branches and Se source bottle and making sure that branches between the
metal sources is always hotter than the Se source. Under these
circumstances Se will always condense into the source not in the
branches. This was achieved by properly sequencing the temperatures
during heat-up and cool-down. Figure 38 shows the relevant temperatures
during a run as a function of time in which the proper sequencing
allowed the Se source to be the coolest point at all times. The efficacy
of this scheme was verified by depositing just Cu right after a CIGS
deposition. No Cu-Se compounds were observed in the resulting film
indicating that no condensation was present in the Se branches.

[]{#__RefHeading___Toc469989169 .anchor}Figure 38. Temperature of the
sources and of the Se branches during a CIGS run where heaters were
properly sequenced during heat-up and cool-down to keep Se source
temperature the lowest. Right hand scale shows temperature

of the Cu, Ga and In sources.

## 3.2 CuInSe~2~-based Devices with Increased V~oc~

## 

The goal of the sub-task on increased V~oc~ is to provide the
understanding of materials and processing requirements to increase the
operating voltage of CuInSe~2~-based solar cells. This can directly
improve the performance of thin film modules. The first part of this
sub-task has focused on the growth of Cu(InAl)Se~2~ films with a
specific objectives to improve the adhesion of Cu(InAl)Se~2~ on glass/Mo
substrates, and determine the effects of increasing relative Al content
and bandgap on material and device properties. A second part of this
work is to develop a fundamental understanding of the diffusion of S
into CuInSe~2~ or Cu(InGa)Se~2~ films which could also be used to
increase V~oc~.

### 3.2.1 Cu(InAl)Se~2~ Solar Cells

Significant effort has been devoted to developing high efficiency
CuInSe~2~-alloy based solar cells with increasing absorber layer bandgap
(E~g~). This is primarily motivated by the expectation that the
resultant increase in voltage and decrease in current under normal
operating conditions can be used to reduce losses in the TCO layers or
to increase interconnect spacing, yielding higher module efficiency. In
addition, the output of wider bandgap cells and modules should have a
reduced temperature coefficient resulting in improved performance under
actual operating conditions. Finally, alloys with E~g~ \> 1.5 eV may
enable these materials to be used as the wide bandgap top cell in a
multijunction solar cell.

Substantial effort continues with CuInSe~2~-based alloys using Ga and/or
S to increase the bandgap but the efficiency of devices with E~g~ \> 1.3
eV is reduced by a degradation of the electronic properties of the
absorber layer leading to losses in FF and V~OC~ [^22] [^23]^.^ We have
previously reported on thin film solar cells using Cu(InAl)Se~2~ [^24]
[^25] [^26] In this alloy system, the bandgap can be continuously varied
from 1.0 to 2.7 eV as x ≡ Al/(In+Al) is increased from 0 to 1.
Cu(InAl)Se~2~ films have lower relative alloy concentration than
Cu(InGa)Se~2~ or CuIn(SeS)~2~ films with comparable bandgap and, as
illustrated in Figure 39, smaller change in lattice parameter.

Cu(InAl)Se~2~ films were deposited by multisource elemental
co-evaporation onto Mo-coated soda lime glass substrates. Poor adhesion
at the Mo/Cu(InAl)Se~2~ interface was identified as a critical issue
Error: Reference source not found and was substantially improved by
incorporating a thin, \~ 5 nm, Ga layer sputtered onto the Mo film prior
to the Cu(InAl)Se~2~ deposition^30^. With better adhesion, films could
be deposited at T~SS~ = 530°C resulting in higher efficiency cells. Most
notably, with Al/(In+Al) = 0.13 and E~g~ = 1.16 eV, a cell with 16.9%
efficiency was demonstrated after application of an anti-reflection
layer ^30^. The Cu(InAl)Se~2~ film structure and device characteristics
were essentially indistinguishable from those with Cu(InGa)Se~2~ that
has Ga/(In+Ga) = 0.25 to give the same bandgap. The primary objective
for exploring Cu(InAl)Se~2~ based solar cells remains the potential for
improved performance with wider bandgap. Two critical issues have
prevented this to date. First, adhesion at the Mo/Cu(InAl)Se~2~
interface is still a problem, even with the Ga seed layer, and becomes
worse with higher Al content. Second, the efficiency decreases for E~g~
≥ 1.4 eV with V~OC~ lower than with the Cu(InGa)Se~2~ based devices. In
this report, the effects of Cu(InAl)Se~2~ deposition with different
flux-temperature-time profiles on film growth and device performance are
presented with emphasis on the film structure, adhesion and device
performance with increasing E~g~. In addition, device behavior as a
function of Cu(InAl)Se~2~ layer bandgap is characterized and compared.

### 3.2.2 Experimental Procedures for Cu(InAl)Se~2~

Cu(InAl)Se~2~ films were deposited by evaporation from Knudsen-type
effusion sources for Cu, In, Al, and Se which provided elemental fluxes
coincident onto radiantly heated soda lime glass/Mo or glass/Mo/Ga
substrates. The sputtered Mo layers were 0.7 µm thick with the sheet
resistance 0.2 Ω/sq. The 5 nm thick Ga layers were deposited directly
onto the Mo in the same sputter system.

Three different flux-temperature-time profiles were used for deposition
of the Cu(InAl)Se~2~ films:

1.  A uniform 1-step process, in which all the fluxes were kept constant
    so there was no Cu-rich composition during the deposition and no
    intentional compositional gradient. The substrate temperature was
    kept constant at T~SS~ = 530°C.

2.  A 2-step process which begins with a Cu-rich Cu(InAl)Se~2~ layer,
    with \[Cu\]\>\[In\]+\[Al\], deposited at T~SS~ = 530°C, followed
    continuously by a layer containing only In, Al and Se at the same
    T~SS~.

3.  A 3-step process which begins with a layer containing only In, Al
    and Se at T~SS~ = 400°C, followed continuously by a layer containing
    only Cu and Se at T~SS~ = 530°C, so that \[Cu\]\>\[In\]+\[Al\], and
    ended by a layer containing only In, Al and Se at T~SS~ = 530°C.

All 3 processes resulted in films with Cu/(In+Al) ≈ 0.7 -- 0.9 and
thickness 2.0 ± 0.5 µm, deposited in \~ 1 hour. These processes were
done with relative Al content varied over the approximate range 0.1 ≤ x
≤ 0.4, corresponding to 1.1 ≤ Eg ≤ 1.5 eV. Both glass/Mo and glass/Mo/Ga
substrates were typically included in a single deposition.

Two additional Cu(InAl)Se~2~ process modifications were made with \~ 5
nm thick In or Cu seed layers evaporated in-situ onto the Mo prior to a
1-step Cu(InAl)Se~2~ deposition. In these cases, there was probably Se
in these layers due to residual Se in the evaporation system. An Al
interlayer was not considered, as Al at the Mo interface seems to cause
the adhesion problem.

The Cu(InAl)Se~2~ films are characterized by x-ray diffraction (XRD),
scanning electron microscopy (SEM), and energy dispersive spectroscopy
(EDS). Optical band gap was determined from the energy at which
interference fringes in the reflection spectrum begin and from quantum
efficiency measurements (QE) on cells. These measurements agree within ±
0.05 e V, with each other and with ellipsometry measurements on selected
samples.

Poor adhesion was determined qualitatively by the delamination of the
Cu(InAl)Se~2~ film during processing, particularly during the CdS
chemical bath deposition. Semi-quantitative adhesion measurements were
also made using an Instron pull tester with a custom sample holder and
coupling apparatus. This test measured the perpendicular breaking point
force needed to remove a stud with area 0.5 cm^2^ which had been epoxied
to the Cu(InAl)Se~2~ surface.

Solar cells were fabricated using the baseline processes developed for
Cu(InGa)Se~2~ devices, with structure glass/Mo/Cu(InAl)Se~2~/CdS/ZnO/ITO
with a Ni-Al grid and no anti-reflection layers. Cells with a total area
of 0.47 cm^2^ were defined by mechanical scribing. The devices were
characterized by current-voltage (J-V) measurements under 100 mW/cm^2^
AM1.5 illumination at 25°C, and by QE and capacitance measurements.

### 3.2.3 Results: Material Characterization

To compare the morphology and grain size with the 1, 2, and 3-step
deposition processes, SEM cross-section images and tapping mode AFM
surface topography scans are shown in Figure 36. The grain structure is
smallest with the uniform process and largest with the 3-step process,
similar to results obtained for Cu(InGa)Se~2~ films [^27] [^28]. The Rms
surface roughness determined from the AFM images were 98 nm for the
1-step process, 138 nm for the 2-step process, and 131 nm for the 3-step
process. This is similar to Cu(InGa)Se~2~ films where a 1-step
deposition yielded smaller grains with lower Rms surface roughness than
a 2-step process Error: Reference source not found

Films deposited with the thin Cu seed layer had grain size and roughness
comparable to the 2-step process and with the In seed layer, the films
were comparable to the 1-step process on either glass/Mo or glass/Mo/Ga
substrates.

+-----------+--------------------------+-------------------------------+
| D         | SEM cross-sectional      | AFM top surface images        |
| eposition | images                   |                               |
|           |                          | (10 x 10 µm)                  |
| Process   | (2.5 µm wide)            |                               |
+-----------+--------------------------+-------------------------------+
| 1-step    | ![](media/image42.bmp)   | ![](media/image43.png)        |
| d         |                          | {width="2.6444444444444444in" |
| eposition |                          | h                             |
|           |                          | eight="2.0902777777777777in"} |
+-----------+--------------------------+-------------------------------+
| 2-step    | ![](media/image44.bmp)   | ![](media/image45.png)        |
| d         |                          | {width="2.6444444444444444in" |
| eposition |                          | h                             |
|           |                          | eight="2.1222222222222222in"} |
+-----------+--------------------------+-------------------------------+
| 3-step    | ![](media/image46.bmp)   | ![](media/image47.png)        |
| d         |                          | {width="2.6444444444444444in" |
| eposition |                          | height="2.154166666666667in"} |
+-----------+--------------------------+-------------------------------+

[]{#__RefHeading___Toc469989170 .anchor}Figure 39. Differences in grain
size and morphology with different

deposition processes as revealed by SEM cross-sectional images of

2.5 µm wide regions, and AFM top surface images of 10 x 10 µm regions.

XRD measurements in a standard θ-2θ configuration showed the films to
vary from a strong (112) to a (220)/(204) preferred orientation. The
degree of orientation of the films along the (112) direction F (112) was
calculated according to the Harris method, by:

\(1\)

where I(hkl) is the measured intensity of the peak (h,k,l) for the film
and Io(hkl) the relative powder intensity. Since no JCPDS file is
available for Cu(InAl)Se~2,~ a CuInSe~2~ standard was used [^29]. For a
film completely (112) oriented, F(112)=1 and for a random oriented film
F(112) = 0.5, since only the (112) and (220)/(204) peaks were included
in the analysis. The 1 and 3-step processes each gave a range of F(112)
with different samples varying from 0.3 -- 0.9, ranging from slightly
(220)/(204) or random to a pronounced (112) orientation. The exception
was the 2-step process and films grown with the thin Cu seed layer. In
these cases, all films had F(112) ≥ 0.95, i.e. they were strongly (112)
oriented, indicating that the Cu seed layer had the same effect as the
Cu-rich growth in the 2-step process.

Adhesion was characterized by whether or not the Cu(InAl)Se~2~ films
adhere during the CBD CdS deposition. This could potentially be
quantified by determining the yield of samples surviving the CBD
process, although to be statistically meaningful, it would require a
large sample set which was not available in this work.

With the 1-step Cu(InAl)Se~2~ deposition process for x ≈ 0.1 using a Ga
seed layer, all films adhere through the bath. As Al content increases
to x \> 0.3 with this process, some samples flake off though most still
adhere. Without the Ga seed layer a few films, \< 30%, still adhered,
but only for x ≈ 0.1. Deposition starting with a Cu seed layer provided
the same improved adhesion as the Ga layer with the 1-step films. The In
seed layer, however, did not provide any improvement in adhesion for
deposition on glass/Mo.

The 2-step process gave no apparent improvement in adhesion with
comparable results to the 1-step process, with or without the Ga seed
layer.

Finally, the 3-step process had much worse adhesion than the 1 or 2-step
depositions, with or without the Ga layer. Films with x = 0.1 frequently
did not adhere through the CdS deposition and with x ≥ 0.3, the films
never adhered and often flaked off, even as they were removed from the
Cu(InAl)Se~2~ deposition system.

Attempts were made to quantitatively determine adhesion using an Instron
pull tester equipped with a custom sample holder and coupling apparatus.
This test measures the perpendicular breaking point force needed to
remove a stud with area 0.5 cm^2^ which had been epoxied to the
Cu(InAl)Se~2~ surface. In most cases, the pulled stud cleanly removed
the Cu(InAl)Se~2~ film from the Mo. A typical plot of perpendicular load
vs. time is shown in Figure 40 yielding a maximum load of 48 N. However,
the force to remove the films varied from 5--30x10^5^ N/m^2^ with little
correlation to the observed adhesion in the CdS bath. This suggests that
there is a large variation in the process, and the poor adhesion through
the CdS process may not be simply related to the mechanical strength of
the Mo/Cu(InAl)Se~2~ interface The exception was the poor adhesion of
films deposited using a 3-step process which begins with an In-Al-Se
layer where the films pulled off with force \<\<5x10^5^ N/m^2^. Films
deposited with this process also rarely adhere in the CdS bath.

![](media/image48.wmf){width="3.1013888888888888in"
height="2.6979166666666665in"}

[]{#__RefHeading___Toc469989171 .anchor}Figure 40. Load needed to remove
a stud epoxied to a

Cu(InAl)Se~2~ surface for measurement of adhesion.

### 3.2.4 Results: Cu(InAl)Se~2~ Solar Cells

J-V parameters for the best cells fabricated from Cu(InAl)Se~2~ using
the three different flux-temperature-time profiles with E~g~ from 1.1 to
1.5 eV are summarized in Table 15. For low bandgap ≤ 1.16 eV, the device
results are essentially independent of the Cu(InAl)Se~2~ deposition
process. For higher Al content, no cell results were obtained with the
3-step process due to the poor adhesion, while the 1 and 2-step
processes gave similar device behavior. As reported previously, the
decrease in efficiency was due to the lack of sufficient increase in
V~OC~ and decrease in FF. It should be noted that the results in Table 1
are the best cells from a number of different depositions. The
differences between J-V parameters with the different processes are most
likely due to statistical variation within relatively small sample
sizes, which vary between the processes.

With E~g~ ≈ 1.1 eV, devices using the Ga seed layer are compared to
devices with the Cu(InAl)Se~2~ deposited directly on glass/Mo in the
same deposition. In each case, the J-V parameters are the same, within
the typical sample-to-sample variations. Poor adhesion without the Ga
layer prevented this comparison with higher bandgap. Films deposited
using the Cu seed layer gave similar device performance to those with
the Ga seed layer. The In seed layer was primarily used only with high
bandgap and the films had poor adhesion.

[]{#__RefHeading___Toc469989109 .anchor}Table 15. Cu(InAl)Se~2~ device
parameters for films with E~g~ = 1.1 to 1.5 deposited with different
processes. With the 3-step process films with E~g~ \> 1.1 did not adhere
during device processing.

+------------+-----------+--------+--------+-----------+-------------+
| E~g~ (eV)  | back      |       | V~OC~  | J~SC~     | FF          |
|            |           |        |        |           |             |
| ± 0.05     | contact   | (%)    | \(V\)  | (         | (%)         |
|            |           |        |        | mA/cm^2^) |             |
+------------+-----------+--------+--------+-----------+-------------+
| *1-step    |           |        |        |           |             |
| process*   |           |        |        |           |             |
+------------+-----------+--------+--------+-----------+-------------+
| 1.10       | Mo/Ga     | 13.0   | 0.525  | 34.8      | 71.4        |
+------------+-----------+--------+--------+-----------+-------------+
| 1.16       | Mo/Ga     | 14.9   | 0.605  | 34.3      | 71.8        |
+------------+-----------+--------+--------+-----------+-------------+
| 1.16       | Mo        | 14.1   | 0.594  | 34.5      | 68.7        |
+------------+-----------+--------+--------+-----------+-------------+
| 1.40       | Mo/Ga     | 13.4   | 0.724  | 26.9      | 69.0        |
+------------+-----------+--------+--------+-----------+-------------+
| 1.50       | Mo/Ga     | 9.9    | 0.747  | 20.1      | 65.8        |
+------------+-----------+--------+--------+-----------+-------------+
| *2-step    |           |        |        |           |             |
| process*   |           |        |        |           |             |
+------------+-----------+--------+--------+-----------+-------------+
| 1.10       | Mo/Ga     | 12.8   | 0.491  | 36.2      | 72.4        |
+------------+-----------+--------+--------+-----------+-------------+
| 1.10       | Mo        | 13.2   | 0.498  | 37.1      | 71.6        |
+------------+-----------+--------+--------+-----------+-------------+
| 1.25       | Mo/Ga     | 13.8   | 0.666  | 29.9      | 69.2        |
+------------+-----------+--------+--------+-----------+-------------+
| 1.40       | Mo/Ga     | 8.7    | 0.688  | 23.1      | 54.4        |
+------------+-----------+--------+--------+-----------+-------------+
| *3-step    |           |        |        |           |             |
| process*   |           |        |        |           |             |
+------------+-----------+--------+--------+-----------+-------------+
| 1.10       | Mo/Ga     | 12.0   | 0.547  | 33.0      | 66.5        |
+------------+-----------+--------+--------+-----------+-------------+

### 3.2.5 Cu(InAl)Se~2~ Device Characterization

## 

The J-V curves of the three devices listed in Table 3.1 using
Cu(InAl)Se~2~ with E~g~ = 1.16, 1.4, and 1.5 eV deposited using the
1-step process on glass/Mo/Ga/substrates are shown in Figure 41. As
previously, these devices show that V~OC~ increases slowly with E~g~, to
a maximum of only \~750 mV, as FF decreases. Also, as E~g~ increases,
there is increasing cross-over of the light and dark J-V curves which
can be seen in Figure 41, suggesting light induced changes in the
absorber layers or at the interface, since the CdS/ZnO/ITO window layers
are nominally identical in all devices.

QE measurements were used to further characterize the effect of
increasing bandgap. The curves are shown in Figure 42 at both the 0V and
--1V bias voltage. The difference between the voltage biases increases
as E~g~ increases as shown by the ratio QE (-1V) / QE (0V) is shown in
Figure 43. For the wider bandgap, this ratio increases dramatically for
longer wavelength light with which an increasing fraction of the
carriers are generated deeper in the absorber layer. This is a clear
indication that the light-generated electrons are not collected
efficiently. The QE voltage dependence is quantified by the quantity
Γ~QE~ which is the ratio of the current obtained by integrating the
product of QE (-1V) times the AM1.5 spectrum to that obtained from QE
(0V). This is also shown in Figure 43.

![](media/image49.wmf){width="3.497916666666667in"
height="2.9145833333333333in"}

[]{#__RefHeading___Toc469989172 .anchor}Figure 41. Light and dark J-V
curves with Cu(InAl)Se~2~

bandgap (a) 1.16 eV, (b) 1.4 eV, and (c) 1.5 eV.

![](media/image50.wmf){width="3.1979166666666665in"
height="2.6645833333333333in"}

[]{#__RefHeading___Toc469989173 .anchor}Figure 42. QE curves of devices
in Figure 38. Solid lines were

measured at 0V and dashed lines at --1V.

![](media/image51.wmf){width="3.497916666666667in"
height="2.9145833333333333in"}

[]{#__RefHeading___Toc469989174 .anchor}Figure 43. Ratio of QE(-1V) /
QE(0V) showing increased

voltage dependence with wider bandgap.

Capacitance-voltage measurements were completed on the devices with 1.16
and 1.4 eV bandgap absorber layers. Measurements were done at frequency
= 100 kHz with the devices at room temperature and used to determine the
depletion width (W) at 0V, assuming that the depletion layer is in the
Cu(InAl)Se~2~ layer. With E~g~ = 1.16 eV, this gave W = 0.3 µm while
with E~g~ = 1.4 eV, W = 0.7 µm. The increase in W with E~g~ appears to
contradict the QE behavior at short wavelength where most of the light
is absorbed in within \~ 0.2 µm, suggesting that many carriers are not
collected, even if generated within the depletion region. More extensive
characterization of the electronic properties of the Cu(InAl)Se~2~
devices using transient photocapacitance spectroscopy and drive level
capacitance profiling was completed in collaboration with the University
of Oregon [^30].

### 3.2.5 Cu(InAl)Se~2~ Discussion And Conclusions

## 

Overall, no benefits for either improved adhesion or device performance
were obtained with the different Cu(InAl)Se~2~ deposition processes.
Both the Ga or Cu seed layers give comparable device behavior to the
SL/Mo with greatly improved yield due to the better adhesion. Thus, two
methods to improve adhesion, with Ga sputtered onto the Mo or Cu
deposited on the Mo as the first step in the Cu(InAl)Se~2~ evaporation
system have been demonstrated. The Cu layer consistently gives (112)
orientation and larger grain size, but the device characteristics are
similar.

Devices with E~g~ ≥ 1.3 eV have decreasing efficiency and the maximum
V~OC~ achieved with Cu(InAl)Se~2~ remains significantly lower than
should be attainable with wide bandgap Cu(InAl)Se~2~ alloy based solar
cells. The recombination current that controls V~OC~ and poor collection
of light-generated minority carriers is likely due to differences in
electronic defects. However, more fundamental characterization will be
needed to determine the nature of the controlling defects.

In general, a number of issues regarding Cu(InAl)Se~2~ thin film solar
cells remain. For example, it has not been determined if the device
structure, identical to that used for most high efficiency Cu(InGa)Se~2~
cells, is optimal for Cu(InAl)Se~2~ absorber layers. In particular,
adhesion and electrical contact might be improved with a different
contact metal. Also, it is not known if device performance, particularly
with wider bandgaps; could be improved by modifying the junction
formation with alternative window layer materials or processing.

The critical issues of poor adhesion and decreasing cell performance at
wide bandgap must be resolved before Cu(InAl)Se~2~ alloys can be
considered a preferred alternative for thin film solar cells. However,
there are a number of issues yet to be explored and, as with the other
CuInSe~2~ based alloys, there is still a lack of fundamental
understanding of important aspects of the materials and devices.

### 3.2.6 CuAlSe~2~ Deposition and Characterization

## 

As part of the effort obtain quaternary Cu(InAl)Se~2~ films by
multi-source physical vapor deposition an effort was undertaken to form
ternary CuAlSe~2~ films with varying Cu/Al ratios to investigate the
phases formed as suggested by the phase diagram in the literature[^31].

Deposition of device quality quaternary CuIn~1-x~Al~x~Se~2~ films as
well as CuAlSe~2~ thin films of varying \[Cu\]/\[Al\] ratios and
thickness (1 -- 4 μm) was carried out in a multi-source physical
deposition system. The films were deposited on glass and molybdenum
coated glass (boroscillicate and soda lime) substrates at three
different substrate temperatures (T~ss~) 350^o^C, 450^o^C and 530^o^C.
The band gap of CuIn~1-x~Al~x~Se~2~ films was varied by controlling x,
the relative amount of Al in the films. All films were analyzed by X-ray
diffraction (XRD) and Energy Dispersive X-ray Spectroscopy (EDS) to
determine their crystal structure and chemical composition,
respectively. Some films were analyzed for their optical properties
using optical spectroscopy or ellipsometry.

It was shown that for Cu-Al-Se films deposited under conditions for
which R~Se~/R~M~ \< 10, where R~Se~ is the effusion rate of Se and R~M~
is the total metal effusion rate (Al + Cu) in mol/min, the resulting
films contain Cu-Al intermetallics (CuAl~2~, CuAl and Cu~2~Al), if the
\[Cu\]/\[Al\] ratios, by EDS, are less than 1.4. The specific
intermetallic observed depends on the \[Cu\]/\[Al\] ratio. Several of
these films also contain Al~2~Se~3~. Composition analysis showed that
these films contain concentrations of oxygen varying from 12 to 60 at. %
as the \[Cu\]/\[Al\] ratio decreases. This oxygen is postulated to come
from the subsequent hydrolysis of Al~x~Se~y~ leading to the formation of
Al(OH)~3~ or Al~2~O~3~ upon exposure to the air. For films whose
composition shows \[Cu\]/\[Al\] ≤ 0.8, the phases identified vary from a
mixture of α-CuAlSe~2~ and Cu~2~Al to a single phase of α-CuAlSe~2~ and
a two-phase mixture of α-CuAlSe~2~ and Cu~2~Se. The oxygen content in
films containing CuAlSe~2~ and Cu~2~Se is at most 6 at %.

For Cu-Al-Se films deposited under conditions for which R~Se~/R~M~ \>
10, with

\[Cu\]/\[Al\] [\<]{.underline}0.8, the resulting films are a mixture of
what seems to be a new structure labeled β-Cu~x~Al~y~Se~2,~ and
α-CuAlSe~2~. In these films β-CuAlSe~2~, is clearly the dominant phase.
These films have optical properties similar to those of α-CuAlSe~2~
obtained from the literature. For films whose composition shows 0.8 \<
\[Cu\]/\[Al\] \< 1.1, the phases identified go through a transition from
β-Cu~x~Al~y~Se~2~ being the dominant phase to αCuAlSe~2~ being the
dominant phase. At \[Cu\]/\[Al\] \~ 1.0, the phases are approximately
equal. As \[Cu\]/\[Al\] \> 1.1, the films are a two-phase mixture of
α-CuAlSe~2~ and Cu~2~Se. The oxygen content in all of these films varies
between 1 - 3 at %. Based on the Cu-Se, Al-Se, Cu-Al binary and the
Cu~2~Se-Al~2~Se~3~ pseudo-binary phase diagrams, a phenomenological film
growth mechanism was developed to explain these results. The results
suggested that the film growth kinetics is controlled by rapid kinetics
of Cu~2~Se formation and also by the difficulty of selenizing Al for Se
delivery such that

R~Se~/R~M~ \< 10.

### 3.2.7 Sulfur Diffusion in Cu(InGa)Se~2~ 

### 

The efficiency of state-of-the-art CuInSe~2~-based solar cells is
limited by recombination through defects in the space charge region,
which leads to high values of the diode current [^32]. Post-deposition
sulfurization of CuInSe~2~-based absorber layers promises to reduce this
recombination process by increasing the bandgap, and therefore the
activation energy for recombination, in the space charge region [^33]
[^34]. This work is aimed at a quantitative understanding of the
transport of S into CuInSe~2~ during post-deposition sulfurization to
enable controlled bandgap modification. Initially, this effort has
focused on diffusion of S into CuInSe~2~, with no Ga, during reaction in
H~2~S.

Sulfur reaction was done in an atmospheric pressure reactor capable of
supplying a mixture of Ar and H~2~S to samples sitting in its hot zone.
Typical reaction conditions were a 1 hour anneal at 575 ºC in 2% H~2~S
with balance Ar. In an effort to eliminate grain boundary and substrate
effects, single-crystal samples grown by various methods were obtained
and sulfurized. These crystals were grown by either vertical gradient
freeze (VGF) or by the traveling heater (TH) method. During TH growth
either In or CuSe were used as solvents. All crystals were polished
after growth. The samples were analyzed with the use of SEM, EDS and
XRD.

VGF and TH(CuSe) grown crystals remained relatively smooth after the
sulfur reaction process, in contrast to TH(In) grown crystals, as seen
in the SEM images in Figure 39. In addition, the TH(In) crystals
incorporated substantially more S than the other crystal types. EDS on
TH(In) crystals after sulfurization resulted in a S/(Se+S) value of 0.73
compared to values of \~0.04 for the other crystal types.
Cross-sectional SEM images of a reacted TH(In) crystal showed a rough
surface layer on top of an undisturbed crystalline bulk. XRD in
Bragg-Brentano geometry featured clearly distinguished CuIn(SeS)~2~
peaks, indicating that this surface layer is polycrystalline. These
differences in sulfurization behavior may result from differences in the
initial Se content of the various crystals.

![](media/image52.jpeg){width="2.0819444444444444in"
height="1.9916666666666667in"} (a) VGF grown crystal

![](media/image53.jpeg){width="2.0819444444444444in"
height="1.9916666666666667in"} (b) TH(CuSe) grown crystal

![](media/image54.jpeg){width="2.1in" height="1.9916666666666667in"} (c)
a TH(In) grown crystal

[]{#__RefHeading___Toc469989175 .anchor}Figure 44. SEM images of
CuInSe~2~ crystals after sulfur reaction.

The TH(In) grown crystal has formed a very rough surface compared

to the other two crystals. Each image shows a 50 x 50 µm area.

To quantitatively characterize the S diffusion, VGF CuInSe~2~ crystal
samples were exposed to a flowing atmosphere of 2% H~2~S in balance Ar
at 575°C. The resulting S distribution in the samples was measured by
Auger Electron Spectroscopy (AES) depth profiling and analyzed on the
basis of the thin-film diffusion model published previously [^35]. This
model assumes a fast chalcogen exchange reaction at the gas-solid
interface and, subsequently, lattice and grain boundary diffusion in the
thin film. It is further assumed that the film structure is conserved
during the sulfurization experiment. The interface reaction quickly
leads to a value very close to 1 for the relative S content y=S/(S+Se)
at the free surface of the film. Consequently, for the analysis of the
diffusion problem at the free film surface a constant concentration
boundary condition of y~s~ =1 is assumed. The only fitting parameters in
this model are the values of the lattice and the grain boundary
diffusion coefficients, D~l~ and D~b~, respectively. Further, it was
assumed that this process could be treated like interdiffusion in a
binary system, despite the fact that four elements (Cu, In, Se, S) were
present. Generally, the lattice diffusion coefficient D~l~ governing
interdiffusion in a binary diffusion couple is a function of composition
y. Its concentration dependence can only be determined from experiments
with one-dimensional diffusion geometry. Therefore, D~l~(y) was obtained
from experiments done on bulk crystalline samples rather than on films.

Three polished VGF crystals were sulfurized for 1, 4 and 9 hr,
respectively. After sulfurization only the one-hour and the nine-hour
sample featured regions that were smooth enough for quantitative
analysis. AES depth profiles were measured on these regions and used to
confirm that the S distribution in the crystals could be described as a
one-dimensional diffusion problem. Profiles created by one-dimensional
diffusion for various diffusion times agree with each other when
presented as concentration versus sputter depth divided by the square
root of the diffusion time [^36]. The measured y profiles are presented
in this way in Figure 45, and the curves agree reasonably well. The very
slight disagreement could be the result of subtle surface imperfections
that affected the diffusion geometry more strongly in the shallower
one-hour profile. This result was crucial since it provided a
self-consistency justification of treating the interdiffusion of S and
Se in CuIn(Se~1-y~S~y~)~2~ as binary interdiffusion.

![](media/image55.png){width="3.9791666666666665in"
height="2.984722222222222in"}

[]{#__RefHeading___Toc469989176 .anchor}Figure 45. AES depth profiles of
the relative S content y measured on crystals sulfurized for 1 hour and
9 hours, respectively. The x-axis shows the sputter depth divided by the
square root of the sulfurization time. The fact that both curves agree
in this graph confirms that the S distribution in the crystals was
shaped by one-dimensional diffusion.

For the determination of D~l~ as a function of y, the better resolved
nine-hour profile was used. In Figure 46, this data is shown after a
Boltzmann-Matano transformation Error: Reference source not found. The
function D~l~(y) has a value between 1 and 2 times 10^-16^ cm^2^s^-1^
for y\>0.4 and increases significantly only for y values smaller than
0.3. For y values below 0.1 the data is very noisy. Figure 46 also
includes a parameterized approximation to the experimental data points
given by:

$D_{1} = 7 - 5\text{.}7\text{cos}\left\lbrack \frac{p}{2}(1 - y)^{2\text{.}7} \right\rbrack\text{10}^{- \text{16}}\text{cm}^{2}s^{- 1}$
(2)

A reverse Boltzmann-Matano transformation of this curve provides a good
fit to the experimental depth profile as shown in Figure 47, clearly
better than the agreement with any of the curves generated with a
constant value of D~l~. Comparison with modeled curves generated with
the use of a constant value for D~l~ shows that the increase towards
lower values of y is significant. Nevertheless, for the investigation of
the S-Se interdiffusion in films, use will be made of the fact that D~l~
turned out to be almost constant over a wide range of concentrations.
For diffusion in a three-dimensional geometry analytical solutions exist
only in the case that D~l~ and D~b~ are constant, and those are
implemented in the thin-film diffusion model. Since the concentration
dependence of D~l~ was found to be limited to a narrow range of
concentrations the model may still be able to approximate the situation
in thin films sufficiently well. Efforts are underway to prepare
sufficiently smooth CuInSe~2~ films to justify the model assumptions
about the diffusion geometry and the measurement procedure.

![](media/image56.png){width="4.499305555555556in"
height="3.3743055555555554in"}

[]{#__RefHeading___Toc469989177 .anchor}Figure 46. D~l~ as a function of
y in CuIn(Se~1-y~S~y~)~2~, calculated, with the help of the
Boltzmann-Matano method, from the AES depth profile of the crystal
sulfurized

for 9 hours. Also shown is a parameterized approximation to this data.

![](media/image57.png){width="4.36875in" height="3.0368055555555555in"}

[]{#__RefHeading___Toc469989178 .anchor}Figure 47. Comparison of the AES
depth profile measured on

the nine-hour sample with various simulated curves.

## 3.3. Fundamental Interface and Material Characterization

### 3.3.1 Emitter Layer Deposition and Junction Characterization

Chemical bath deposited ZnS buffer layers have been used to form the
highest efficiency Cu(InGa)Se~2~-based solar cells [^37] which contain
no CdS layer. However, it is difficult to explain how such high
efficiencies can be obtained with a conduction band offset of over 1.2
eV [^38] between the Cu(InGa)Se~2~ and ZnS. A possible explanation for
the device operation is that oxygen must be incorporated in the CBD-ZnS
to form Zn(OS) to decrease the conduction band offset between
Cu(InGa)Se~2~ and Zn(OS). The conduction band of ZnO is about 0.4 eV
lower than Cu(InGa)Se~2,~, and the change in conduction band minimum is
1.6 eV for a change in bandgap of only 0.4 eV. In practice, the high
efficiency devices were annealed in air after the CBD process \[1\].
Typically, this anneal is done at 150°C for 30 -- 60 min.

Experiments were performed to determine the effect of the annealing on
the characteristics of CBD-ZnS films and Cu(InGa)Se~2~/ZnS devices. A
ZnS film was deposited on quartz substrates using the CBD method
developed previously at IEC [^39]. This uses a bath with 25 mM ZnSO~4~,
0.27M thiourea; 2.9M NH~3~H~2~O and temperature = 80°C. The films were
sequentially annealed at 150°C or 200°C for times from 30 -- 360 min and
characterized by optical transmission (T) and reflection (R)
measurements. The quartz substrate enables transmission measurements up
to 5 eV while a soda lime glass substrate will have significant
absorption for hν \> 3.5 eV. Also, glancing angle XRD measurements are
being completed to determine if there is a shift in lattice parameter
due to oxygen incorporation in the films. Finally, devices are also
being fabricated with CBD ZnS layers deposited on Cu(InGa)Se~2~ and
annealed for comparable conditions.

![](media/image58.wmf){width="3.497916666666667in"
height="3.0555555555555554in"}

[]{#__RefHeading___Toc469989179 .anchor}Figure 48. Normalized
transmissions of CBD ZnS film with various air

annealing and a sputtered ZnO film. Labels correspond to Table 16.

[]{#__RefHeading___Toc469989110 .anchor}Table 16. Energy Bandgap of CBD
ZnS and sputtered ZnO

  -------- -------------------- --------------- --------------- --------------
  Sample                        Anneal Temp\    Anneal time\    Eg\
                                (°C)            (min)           (eV)

  1        CBD ZnS              As-deposited                    3.78

  2        CBD ZnS              150             30              3.65

  3        CBD ZnS              200             30              3.59

  4        CBD ZnS              200             150             3.49

  5        CBD ZnS              200             360             3.47

  6        Sputtered ZnO        As-deposited                    3.25
  -------- -------------------- --------------- --------------- --------------

A ZnS film of 170 nm thick (measured by a Dektak profilometer) was
deposited from four sequential baths. Layers used in devices were only
40 -- 50 nm \[45\], but thicker layers were desired for film
characterization. Figure 48 shows the normalized transmission, T/(1-R),
of the as-deposited film and after several anneals at increasing time
and temperature. T/(1-R) of a sputtered ZnO film with thickness 130 nm
is also shown as a comparison. The optical bandgap of the films was
calculated from the plot of (αE)^2^ versus the incident energy, as shown
in Figure 49 and the bandgaps are listed in Table 16. While this
analysis assumes a uniform bandgap, the value of Eg will be that of the
lowest bandgap part of a graded film. The minimum bandgap of the CBD ZnS
decreased from the initial 3.78 eV to 3.47 eV with air annealing at
200°C. Further decrease down to the 3.25 eV of ZnO was not achieved with
further annealing and might require higher temperature. These results
suggest that ZnS(O) alloy films were formed and devices with these
layers can be used to study the effect of the band alignment at the
Cu(InGa)Se~2~/Zn(SO) interface. XRD and device results will be completed
and shown in the next report.

![](media/image59.wmf){width="3.497916666666667in"
height="3.0479166666666666in"}

[]{#__RefHeading___Toc469989180 .anchor}Figure 49. Bandgap determination
from (αE)^2^ versus E for the

ZnS and ZnO samples in Figure 48 with labels corresponding to Table 16.

To determine if the anneal also gave a lattice constant shift consistent
with the alloying of the film, glancing incident angle XRD (GIXRD)
measurements were carried out with 0.5^o^ incident angle using Cu Ka
radiation at 40 KV, 40 mA. Figure 50 shows the GIXRD patterns of 1) as
deposited ZnS/Mo; 2) air annealed ZnS/Mo at 200^o^C for 360 min; 3) air
annealed ZnS/quartz at 200^o^C for 360 min and 4) sputtered ZnO/quartz
for comparison. There was no apparent change in the (111) peak position
with air annealing in the XRD pattern of the CBD-ZnS films with no
indication of alloying despite the change in optical transmission..

Cell results are summarized in Table 17 for Cu(InGa)Se~2~/ZnS cells with
40 or 80 nm thick ZnS (single or double dip coatings) which received
different anneals. In general there is no clear correlation between the
anneal time and the device performance and, in fact, the best cell was
obtained with no post-deposition anneal. The cells have fill factors as
low as 20 % and J~SC~ as low as 5 mA/cm^2^ due to the appearance of a
inflection in the power quadrant J-V curve in some cases. This is an
indication of a collection barrier, probably at the
Cu(InGa)Se~2~/ZnS(S~1-x~O~x~) interface. The lack of correlation with
the annealing time and temperature is not understood.

![](media/image60.wmf){width="3.497916666666667in"
height="3.0479166666666666in"}

[]{#__RefHeading___Toc469989181 .anchor}Figure 50. GXRD patterns: 1) as
deposited ZnS/Mo; 2) air annealed

ZnS/Mo at 200^o^C for 360 min.; 3) air annealed ZnS/quartz at 200^o^C
for

360 min. and 4) as sputtered ZnO/quartz.

[]{#__RefHeading___Toc469989111 .anchor}Table 17. Summary of device
results with different annealing times for Cu(InGa)Se~2~ cells with CBD
ZnS. Results include the best cell and average of 6 cells on each piece.

+---------+--------+-------+------+-----+-------+-------+------+----+
| Sample  | ZnS    | post  |      | J-V | V~OC~ | J~SC~ | FF   | e  |
|         |        | -dep. |      |     |       |       |      | ff |
|         |        | a     |      |     |       |       |      |    |
|         |        | nneal |      |     |       |       |      |    |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         | thi    | Temp  | Time | t   | (mV)  | (mA/c | (%)  | (  |
|         | ckness | (°C)  | (    | est |       | m^2^) |      | %) |
|         |        |       | min) |     |       |       |      |    |
|         | (nm)   |       |      |     |       |       |      |    |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~40   | None  |      | B   | 602   | 30.8  | 64   | 11 |
| 3726.12 |        |       |      | est |       |       |      | .8 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 604   | 27.7  | 43   | 7  |
|         |        |       |      |     |       |       |      | .4 |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~40   | 150   | 30   | B   | 589   | 22.2  | 29   | 3  |
| 3726.23 |        |       |      | est |       |       |      | .8 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 520   | 14.0  | 27   | 2  |
|         |        |       |      |     |       |       |      | .1 |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~40   | 200   | 30   | B   | 606   | 24.6  | 59   | 8  |
| 3726.33 |        |       |      | est |       |       |      | .8 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 591   | 18.2  | 48   | 5  |
|         |        |       |      |     |       |       |      | .3 |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~40   | 200   | 150  | B   | 594   | 20.9  | 30   | 3  |
| 3725.21 |        |       |      | est |       |       |      | .7 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 547   | 12.6  | 26   | 1  |
|         |        |       |      |     |       |       |      | .9 |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~80   | None  |      | B   | 514   | 9.7   | 21   | 1  |
| 3725.12 |        |       |      | est |       |       |      | .1 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 458   | 5.0   | 22   | 0  |
|         |        |       |      |     |       |       |      | .5 |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~80   | 150   | 30   | B   | 601   | 11.8  | 43   | 3  |
| 3725.22 |        |       |      | est |       |       |      | .1 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 417   | 7.7   | 31   | 1  |
|         |        |       |      |     |       |       |      | .4 |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~80   | 200   | 30   | B   | 578   | 31.1  | 49   | 8  |
| 3725.23 |        |       |      | est |       |       |      | .9 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 483   | 20.3  | 33   | 3  |
|         |        |       |      |     |       |       |      | .9 |
+---------+--------+-------+------+-----+-------+-------+------+----+
| 3       | \~80   | 200   | 150  | B   | 548   | 25.3  | 29   | 4  |
| 3725.31 |        |       |      | est |       |       |      | .1 |
+---------+--------+-------+------+-----+-------+-------+------+----+
|         |        |       |      | Avg | 552   | 17.3  | 25   | 2  |
|         |        |       |      |     |       |       |      | .5 |
+---------+--------+-------+------+-----+-------+-------+------+----+

### 3.3.2 Alternative CdS Deposition

Chemical surface deposition (CSD) of CdS films is described in section
2.4 of this report. In this process, heat from the substrate provides
the energy needed for reaction at the surface, resulting in
heterogeneous deposition. Deposition from an aqueous solution containing
cadmium species, thiourea and ammonia yields dense conformal CdS films,
with \> 90% utilization of Cd species.

Cu(InGa)Se~2~/CdS devices were fabricated using CSD CdS on Cu(InGa)Se~2~
films [^40]. The Cu(InGa)Se~2~ was deposited by elemental evaporation at
550°C and was 2 μm thick with Ga/(Ga+In) = 0.15. The CdS coating was
applied to three structures consisting of soda lime
glass/Mo/Cu(InGa)Se~2~ from a single deposition. Sample 1 was a control
sample with 40 nm CdS deposited by the baseline CBD process. Sample 2
had the CdS deposited by CSD using the following volumetric proportions:
2.8 ml of 0.015 M CdSO~4~ + 2.2 ml of 1.5 M CS(NH~2~)2 + 2.2 ml of 14 M
NH~4~OH + 15 ml of H~2~O. The sample was pre-heated for 10 minutes to
80°C. Sample 3 had the CdS deposited by chemical surface deposition with
two coats, 3 min each of the same solution as sample 2. Solar cell
devices were completed with the sputter deposition of an Al-doped ZnO
layer which was 0.5 µm thick and had sheet resistance of 20 Ω/square,
followed by a Ni/Al collector grid.

Current-voltage (J-V) parameters summarized in Table 18 show that sample
3 had comparable V~oc~ and FF to the control cell. These results
demonstrate that good cell efficiency can be achieved with the high
utilization CdS deposition process. Further optimization of the CSD
process for Cu(InGa)Se~2~ solar cells will focus on the surface
condition prior to deposition and optimization of the CDS conditions for
maximum cell performance.

[]{#__RefHeading___Toc469989112 .anchor}Table 18. Cu(InGa)Se~2~/CdS
device results using chemical surface deposition for the CdS.

+---------+--------+--------+------------+---------+-----------------+
| Sample  | CdS    | V~OC~  | J~SC~      | FF      | Eff             |
| \#      |        |        |            |         |                 |
|         | Dep    | (      | (mA/cm^2^) | (%)     | (%)             |
|         |        | Volts) |            |         |                 |
+---------+--------+--------+------------+---------+-----------------+
| 1       | CBD    | 630    | 32.3s      | 75.6    | 15.4            |
+---------+--------+--------+------------+---------+-----------------+
| 2       | CSD    | 630    | 29.6       | 64.8    | 12.2            |
|         | single |        |            |         |                 |
+---------+--------+--------+------------+---------+-----------------+
| 3       | CSD    | 620    | 28.4       | 75.1    | 13.2            |
|         | double |        |            |         |                 |
+---------+--------+--------+------------+---------+-----------------+

### 3.3.3 Cu(InGa)Se~2~ Optical Characterization

Polycrystalline CuIn~1‑x~Ga~x~Se~2~ thin films are promising materials
for photovoltaic applications. Precise knowledge of the optical
properties of these films is an important factor in characterizing and
modeling these solar cell devices. There is a significant amount of
information available on band structure based on theoretical

calculations [^41] [^42] [^43].

The CuIn~1‑x~Ga~x~Se~2~ alloy system has been investigated on single or
polycrystalline ingots using spectroscopic ellipsometer (SE) [^44] [^45]
[^46]. However, these studies do not address the possible effects of
grain size and polycrystalline texture on the optical properties. SE
measurements on CuIn~1‑x~Ga~x~Se~2~ films reported in the literature
could not be analyzed due to the high film surface roughness and
electronic transitions as a function of the Ga content were not
determined [^47] [^48]. Optical transmission, reflection and absorption
studies on thin film CuIn~1‑x~Ga~x~Se~2~ have generally been limited to
the fundamental absorption edge [^49] [^50] or did not have sufficient
sensitivity to separate the effects of spin orbit and crystal field
interactions near the fundamental band gap. The spin-orbit and crystal
field splittings have been determined by photoluminescence measurements
of CuIn~1‑x~Ga~x~Se~2~ epitaxial films [^51]. However, there is no
optical data available in the literature for polycrystalline
CuIn~1‑x~Ga~x~Se~2~ thin films showing the influence of Ga content on
various electronic transitions over a large spectral range. In this
work, the effect of Ga on the optical properties and electronic
transitions of CuIn~1‑x~Ga~x~Se~2~ polycrystalline thin films is
determined.

In general, it is difficult to model and uniquely determine the optical
constants from ellipsometric data on materials with high surface
roughnessError: Reference source not found^.^ This report describes a
method to circumvent the problem associated with the surface roughness
by peeling the films from the substrate and performing measurements on
the smooth exposed film surface. In addition, a detailed spectroscopic
ellipsometer measurement and modeling routine for CuIn~1‑x~Ga~x~Se~2~
polycrystalline thin films is presented. Accurate optical functions for
the energy range from 0.75 to 4.6 eV are presented for films with x ≡
Ga/(In+Ga) from 0 to 1.

### 3.3.4 Ellipsometry Measurement Strategy 

AFM measurements showed that the evaporated CuIn~1‑x~Ga~x~Se~2~ films
have surface root mean square (R~MS~) roughness on the order of 50 to
100 nm. It is almost impossible to carry out useful SE measurements
because such high surface roughness will not allow the optical functions
to be uniquely determined. Mechanical polishing can be used to smooth
the films, but it is difficult to ensure that the grit is completely
removed from the surface and grain boundaries after polishing. Chemical
polishing could be another method, but etching may modify the film's
surface composition to which SE measurements are highly sensitive.
Measurement through the substrate (reverse ellipsometry) is another
option, provided the films were deposited on optically transparent (e.g.
glass without Mo coating) substrates. However, CuIn~1‑x~Ga~x~Se~2~
deposited on a glass substrate may have different properties than
material grown on a glass/Mo substrate. For example, films on a glass
substrate were shown to have a much stronger (112) preferred orientation
[^52]. Also, with a glass substrate, very thick (\~5 mm) glass is
required to avoid front reflection from the glass affecting the
measurement.

An alternative way to circumvent the surface roughness is to perform the
SE measurements on the reverse side of the film after peeling it from
the substrate. To prepare a sample, an optically thick (\>0.1μm) Mo film
was first deposited on the CuIn~1‑x~Ga~x~Se~2~ film. Then an aluminum
substrate was glued to the film using a fast curing epoxy. Because of
the optically thick Mo, optical functions of the epoxy are not required
for the optical modeling. After curing, the Mo/Glass substrate was
pulled apart from the CuIn~1‑x~Ga~x~Se~2~ film, leaving the smooth and
oxide free back surface of the CuIn~1‑x~Ga~x~Se~2~ film exposed for the
SE measurement. In the present configuration, the original rough surface
is now on the reverse side, where only higher wavelength probing light
can reach. At higher wavelength, the effect of surface roughness is
negligible due to the reduced scattering. For the transparent region,
the ellipsometry data has contributions from the interface layer and the
Mo film. These contributions were modeled with an effective medium
layer, described in a later section. In addition, CuIn~1‑x~Ga~x~Se~2~
films were sufficiently thick so that optical constants in the
transparent region could be determined unambiguously. Measurements and
fits were carried out at three different incident angles to improve the
confidence limit of the data fit in the absorbing region.

SE measurements were carried out using a J.A. Woollam variable angle
spectroscopic ellipsometer (VASE). The VASE is a rotating analyzer
ellipsometer equipped with an auto-retarder, which is useful in
measuring the depolarization caused by surface roughness and thickness
non-uniformity. Measurements were carried out at room temperature
immediately after peeling with the sample under Ar flow to reduce
possible surface contamination. First, VASE scans were done with varying
angle to determine the pseudo-Brewster angle of the multilayer
structure. For this, scans were performed at two wavelengths (500 and
1500 nm) by varying the angle of incidence from 50° to 85° in 1° steps.
The angle corresponding to ∆=90 and the minimum in ψ provides the
pseudo-Brewster angle for that sample [^53]. Here ∆ and ψ are the
ellipsometry parameters corresponding to phase and amplitude changes
induced by the reflection. Then measurements were done at this angle and
two additional angles selected by an empirical rule; pseudo-Brewster
angle ± 5° [^54]. VASE measurements were carried out between 0.75 eV to
2 eV with an energy step of 5 meV and between 2 eV to 4.6 eV with a 12.5
meV step. A small step size in the transparent region was selected
specifically to resolve the closely spaced fundamental transitions and
the interference fringes. The data acquisition and modeling were
performed with the help of WVASE 32 software supplied by J. A. Woollam
Company.

### 3.3.5 Optical Modeling

Chalcopyrite CuInSe~2~ single crystals exhibit uniaxial anisotropy in
their optical

properties [^55] displaying ordinary optical properties (electric field
perpendicular to c-axis) and extraordinary optical properties (electric
field parallel to the c-axis). For polycrystalline material, the
anisotropy is a strong function of the crystallographic texture. If all
the crystals have the same texture, the anisotropy of the
polycrystalline material would be similar to that of a single crystal.
However, if the polycrystalline material has the same probability of
occurrence for all textures or is randomly structured the optical
properties of the material would be isotropic. The effective dielectric
function of the polycrystalline material was calculated as a tensorial
average considering a medium composed of uniaxial crystals oriented in
random directions [^56]. A Bruggeman effective medium calculation shows
that the pseudo-dielectric function of the randomly oriented
polycrystalline material resembles the ordinary component of the single
crystal material [^57]

Extracting useful information from ellipsometric data begins with the
construction of an optical model for the peeled CuIn~1‑x~Ga~x~Se~2~ film
on Mo substrate. Figure 51 shows the structure of the
CuIn~1‑x~Ga~x~Se~2~ sample and the optical model used for the SE
measurement and analysis. The optical model consists of: 1) surface
roughness on the peeled layer; 2) the CuIn~1‑x~Ga~x~Se~2~ film; 3) an
interface layer (surface layer before peeling); and 4) the Mo layer. It
had to be assumed that the optical constants of the Mo film were
different from bulk optical properties. Therefore, the Mo optical
constants were obtained from measurements on Mo/ CuIn~1‑x~Ga~x~Se~2~
films after depositing the \>0.1 μm thick Mo film on the
CuIn~1‑x~Ga~x~Se~2~ and prior to attaching the Al substrate. Optical
constants of the Mo layer were then extracted from a 2-layer model
consisting of a Mo bulk layer and a surface roughness layer, where the
surface roughness layer is modeled using a Bruggeman effective medium
approximation (EMA) consisting of Mo and a void fraction. Here it is
assumed that the surface features are hemispherical in shape and the
corresponding screening parameter is 0.33. Since the Mo film was
optically thick, the effect of the CuIn~1‑x~Ga~x~Se~2~ substrate was not
considered for this analysis.

The complex dielectric response, ε = ε~1~+iε~2~, is the manifestation of
different interband electronic transitions resulting from the applied
electric field. The relation between the energy-band structure and ε~2~
can be given by Error: Reference source not found:

$e_{2}(E) = \frac{4e^{2}h^{2}}{\text{pm}^{2}E^{2}}\int_{}^{}{\text{dk}\left| P_{\text{cv}}(k) \right|}^{2}d\left\lbrack E_{c}(k) - E_{v}(k) - E \right\rbrack$
(3)

where μ is the combined density of states (DOS) mass; the Dirac δ
function represents the spectral joint DOS between the valance
\[E~v~(k)\] and conduction \[E~c~(k)\] band states differing by energy
$E = hw$ of the incident light; P~cv~(k) is the momentum-matrix element
between the valence and conduction-band states; and the integration is
performed over the first Brillouin zone (BZ). Critical points (CP)
correspond to the energy at which the joint DOS shows strong variations
as a function of energy. In three dimensions, depending on the
analytical behavior of the joint DOS, there are four types of CP's,
namely M~0~, M~1~, M~2~ and M~3~ present. The energy-band diagram for a
generic I-III-VI~2~ chalcopyrite structure without spin orbit
interaction is shown in Ref. 45. The chalcopyrite band diagram is
derived from the zinc blend (ZB) structure because of the similarity in
physical properties due to its tetragonal structure [^58].

![](media/image61.wmf){width="3.3645833333333335in"
height="4.847222222222222in"}

[]{#__RefHeading___Toc469989182 .anchor}Figure 51. Sample structure and
optical model

used for the SE measurement and analysis.

The valance band maximum at the BZ center is a triply degenerate state.
The perturbation due to the tetragonal crystal field resulting from the
change in symmetry from ZB to chalcopyrite leads to the non-degenerate
states Γ~4~ and doubly degenerate Γ~5~. Theoretically the crystal field
splitting ∆~cf~ can be calculated using the compression formula [^59]

$D_{\text{cf}}\text{=-}b\left\lbrack 1 - \frac{c}{2a} \right\rbrack$ (4)

where c and a are the lattice constants and b is the deformation
potential with a typical magnitude of \~1eV. Based on this relation the
∆~cf~ changes from positive for CuInSe~2~ (c/2a=1.004) to negative for
CuGaSe~2~ (c/2a=0.9825). The doubly degenerate Γ~5~ is further split by
the spin-orbit interaction.

The fundamental transitions from the valence band maximum to the
conduction band minimum at the Γ point are schematically shown in
Figures 52 (a) and (b). Vertical solid and dashed lines indicate
transitions allowed in perpendicular or parallel electric fields
respectively. For a CuInSe~2~-like CH structure (Figure 38a) ∆~cf~ is
positive and Γ~4v~ lies below Γ~5v.~ For a CuGaSe~2~-like CH structure
(Figure 38b) ∆~cf~ is negative and Γ~5v~ lies below Γ~4v.~ The
transitions are conventionally labeled as E~0~(A), E~0~(B) and E~0~(C)
in the order of increasing energy [^60] The ordering of the energy
levels is determined from the polarization- intensity ratios. The
polarization- intensity ratio for the transitions from the
$\Gamma_{7v}^{4}$and $\Gamma_{7v}^{5}$valance bands to the
$\Gamma_{6c}^{1}$conduction band are given by Error: Reference source
not found

$\frac{I_{||}}{I_{\hat{}}} = \left\lbrack 2 - \frac{3DE}{D_{\text{so}}} \right\rbrack^{2}$
(5)

where ∆E is the separation from the totally perpendicularly polarized
$\Gamma_{6v}^{5}$→ $\Gamma_{6c}^{1}$ transition. The calculated
intensity ratios for $\Gamma_{7v}^{4}$,$\Gamma_{7v}^{5}$and
$\Gamma_{6v}^{5}$ transitions in CuGaSe~2~ crystal are 5:1, 1:4 and 1:1.
Experimental determination of these intensity ratios would require
measurements on single crystals with the c axis perpendicular and
parallel to the incident light. In our analysis we used the analogy of
the single crystal measurements and theoretical calculations to identify
the different energy levels.~.~

![](media/image62.wmf){width="2.8006944444444444in"
height="2.147222222222222in"}
![](media/image63.wmf){width="2.8006944444444444in"
height="2.167361111111111in"}

[]{#__RefHeading___Toc469989183 .anchor}Figure 52. Schematic diagram
showing how the crystal field

and spin orbit interaction leads to formation of the non-degenerate
energy

levels E~0~(A), E~0~(B) and E~0~(C) in (a) CuInSe~2~ and (b) CuGaSe~2~
chalcopyrite materials. Dashed and solid arrows represent optical
transitions allowed

in E\|\|c and E⊥c, respectively. Notation is from Ref. Error: Reference
source not found.

Experimentally, the transition energies can be obtained from CP's in the
~2~ spectrum of the material, and are best resolved in their numerical
second order derivative. The unbiased values of spin orbit splitting
(i.e. spin-orbit splitting in the absence of crystal field splitting)
and crystal field splitting (i.e. crystal field splitting in the absence
of spin-orbit splitting) were determined using a quasi-cubic model
[^61]:

$D_{\text{so}} = \frac{E_{1} + E_{2}}{2} + \left\lbrack \frac{3}{2}\left( \frac{E_{2} - E_{1}}{2} \right)^{2} - \frac{1}{2}\left( \frac{E_{2} + E_{1}}{2} \right)^{2} \right\rbrack^{\frac{1}{2}}$
(6)

$D_{\text{cf}} = \frac{E_{1} + E_{2}}{2} - \left\lbrack \frac{3}{2}\left( \frac{E_{2} - E_{1}}{2} \right)^{2} - \frac{1}{2}\left( \frac{E_{2} + E_{1}}{2} \right)^{2} \right\rbrack^{\frac{1}{2}}$
(7)

where E~1~ = $\Gamma_{7v}^{4}$valance band energy --
$\Gamma_{6v}^{5}$valance band energy and E~2~ = $\Gamma_{7v}^{5}$valance
band energy -- $\Gamma_{6v}^{5}$valance band energy.

Both the surface and interface layers in the optical model for the
peeled film were analyzed with an EMA model. The RMS surface roughness
measured by AFM on the CuIn~1‑x~Ga~x~Se~2~ film before Mo deposition
provided an estimate of the interface layer thickness. The thickness of
the CuIn~1‑x~Ga~x~Se~2~ film was obtained from fitting the data in the
transparent region of the spectrum using a Cauchy dispersion relation.
Then the results for *n* in the transparent region obtained from Cauchy
dispersion relation were used to perform a point-by-point fit starting
from the transparent region, to extract approximate values of the
optical functions. The error in the point-by-point fit, measured by a
90% confidence limit, and was determined to check the quality of the *n*
and *k* data. The optical functions thus obtained are called initial
spectra and are close to the bulk dielectric response of the bulk film,
free of substrate, interface, and surface roughness effects.

The initial spectrum was fit with a General Oscillator model, which uses
a linear summation of Kramers-Kronig (KK) consistent oscillators to
describe the various optical transitions. The fundamental transitions
E~0~(A), E~0~(B), and E~0~(C) in the chalcopyrite crystal have
three-dimensional M~0~ CP character \[44\]. Three parametric Gaussian
Broadened Polynomial Superposition (GBPS) semiconductor oscillators
[^62] are used for simulating these transitions. The CP's at higher
energies due to transitions at Γ, N and T points in the BZ were
simulated using several harmonic Lorentzean and Gaussian oscillators.
The strength, broadening and positions for these oscillators were
adjusted to match the initial spectrum. In the next procedure, the
expected ellipsometric parameters ψ and ∆ of the multilayer structure
were generated using this parametric dispersion layer in the optical
model. The generated data are compared with the experimental data and
fit to reduce the error using the Marquardt-Levenberg algorithm [^63].
In this procedure, the thickness and refractive index of the film are
slightly adjusted. Also, non-ideal features like thickness
non-uniformity are introduced into the model to fit the depolarization
measured using the autoretarder. The root mean square error (MSE) is
used to quantify the difference between the experimental and predicted
data. The MSE is calculated by:

$\text{MSE} = \left\{ \frac{1}{2N - M}\sum_{i = 1}^{N}\left\lbrack \left( \frac{Y_{i^{\text{mod}}} - Y_{i^{\text{exp}}}}{s_{Yj}^{\text{exp}}} \right)^{2} + \left( \frac{D_{i^{\text{mod}}} - D_{i^{\text{exp}}}}{s_{Dj}^{\text{exp}}} \right)^{2} \right\rbrack \right\}^{1/2}$
(8)

where $\sigma_{\Psi j,\Delta j}^{\text{exp}}$ is the standard deviation
in the measured ellipsometry parameters Ψ and ∆,
$\Psi_{i}^{\text{mod}}$and $\Delta_{i}^{\text{mod}}$are the model
generated ellipsometry parameters, $\Psi_{i}^{\text{exp}}$and
$\Delta_{i}^{\text{exp}}$ are the measured ellipsometry parameters , N
is the number of measured parameters, and M is the number of fitted
parameters. By using the experimental standard deviation as the
weighting parameter in the fit, the contribution due to noise in the MSE
is reduced significantly. The sequence of obtaining the point--by point
fit, adjusting the oscillator strength and then fitting to reduce the
MSE was repeated several times until a minimum value of MSE was obtained
without significant parameter correlation. A good fit is obtained when
MSE is close to unity.

### 3.3.6 Cu(InGa)Se~2~ Ellipsometry: Results And Discussion

CuIn~1‑x~Ga~x~Se~2~ films peeled off of the glass/Mo substrate without
significant cracking in the film. EDS measurements on the Mo substrate
after removing the CuIn~1‑x~Ga~x~Se~2~ show the presence of Se. SE
measurements of the Mo substrate indicate the presence of an \~8 nm
thick MoSe~2~ film. This MoSe~2~ layer weakens the interface between the
CuIn~1‑x~Ga~x~Se~2~ film and Mo/Glass substrate allowing for clean
separation. The peeled CuIn~1‑x~Ga~x~Se~2~ film surface is free of
MoSe~2~ and x-ray photoelectron spectroscopy (XPS) measurement shows
that there is no Mo present on the CuIn~1‑x~Ga~x~Se~2~ film.

XRD and EDS measurements show that the CuIn~1‑x~Ga~x~Se~2~ films are
single phase throughout the whole composition range. Table 19 shows the
composition of all CuIn~1‑x~Ga~x~Se~2~ films used in this work, as
determined from EDS measurements. The orientation of the crystals with
respect to the direction normal to the substrate plane was compared to a
random powder orientation using the method of Harris for polycrystalline
fiber texture analysis: [^64]. In this analysis, the orientation
parameter P~hkl~ is given by:

$P_{\text{hkl}} = n\frac{I_{\text{hkl}}}{I_{\text{hkl}}^{r}}\left\lbrack \sum_{\text{hkl} = 1}^{n}\frac{I_{\text{hkl}}}{I_{\text{hkl}}^{r}} \right\rbrack^{- 1}$
(9)

where I~hkl~ is the intensity of the reflection from the hkl plane of
the film, $I_{\text{hkl}}^{r}$ is the intensity of the reflection from
the hkl plane of a reference powder sample with the same phase and n is
the number of peaks included in the analysis. The films in this work
were all found to have a (112) preferred orientation. A film with all
the crystals oriented in the (112) direction normal to the substrate
plane will have P~112~ = n, a film with random grain orientation will
have P~112~ = 1, and a film with grains oriented in a direction other
than (112) will have P~112~ \< 1. In this work, 5 peaks were used for
the orientation analysis and the films have orientation parameter P~112~
in the range 1.8 -- 2.8, as listed in Table 20. This indicates that the
films are neither randomly nor perfectly oriented.

[]{#__RefHeading___Toc469989113 .anchor}Table 19. EDS data of
Cu(In~1-x~Ga~x~)Se~2~ films used for the SE analysis.

  --------- --------- -------- --------- -------- ------------ ------------
  Sample    Cu (%)    In (%)   Ga (%)    Se (%)   Cu/(In+Ga)   Ga/(In+Ga)

  A         25.8      25.2     0         49.0     1.02         0.00

  B         24.8      18.3     8.3       48.7     0.93         0.31

  C         24.0      14.9     12.1      49.0     0.89         0.45

  D         24.6      8.9      17.5      48.9     0.93         0.66

  E         23.7      0        27.3      49.0     0.87         1.00
  --------- --------- -------- --------- -------- ------------ ------------

[]{#__RefHeading___Toc469989114 .anchor}Table 20. Comparison of
orientation parameter, roughness data from

AFM and SE and fit index.

+------------+-----------------+------------------+------------------+
| Ga/(In+Ga) | Orientation     | Roughness from   | Roughness from   |
|            | parameter       |                  |                  |
|            | (P~hkl~)        | AFM (nm)         | SE (nm)          |
+------------+-----------------+------------------+------------------+
| 0          | 2.2             | 5.6              | 3.5              |
+------------+-----------------+------------------+------------------+
| 0.31       | 1.8             | 3.7              | 5.2              |
+------------+-----------------+------------------+------------------+
| 0.45       | 2.7             | 3.4              | 6.1              |
+------------+-----------------+------------------+------------------+
| 0.66       | 2.8             | 14.6             | 8.8              |
+------------+-----------------+------------------+------------------+
| 1.00       | 2.8             | 3.1              | 3.2              |
+------------+-----------------+------------------+------------------+

Figure 53 shows the pseudo-dielectric function measured for CuInSe~2~
film along with the ordinary and extraordinary results for a single
crystal reported in the literature ^52^ . The optical constants of the
CuInSe~2~ film are closer to the ordinary component of the single
crystal optical constants. Therefore, the (112) orientation of the films
does not cause a measurable change in the optical constants. To confirm
this, measurements were performed on these films to determine the
off-diagonal elements in the Jones matrix [^65]. The data was acquired
with the polarizer set at --60°, -30°, 0, 30° and 60°. These anisotropy
measurements showed that the off-diagonal elements were close to zero
confirming that the optical properties are due to the ordinary
component. Also, the magnitude of the ε~2~ spectrum is similar to the
single crystal suggesting that the films are optically dense with
insignificant surface roughness and/or oxide layer. Voids, oxide layers
or roughness can reduce the dielectric response of the material
resulting in a decrease in this ε~2~ spectrum.

![](media/image64.wmf){width="3.4027777777777777in"
height="2.513888888888889in"}

[]{#__RefHeading___Toc469989184 .anchor}Figure 53. The ε~2~ spectrum for
the CuInSe~2~ thin film along with the

ordinary (O) and extra-ordinary (E) ε~2~ spectra for a CuInSe~2~ single
crystal

from Reference 45. The film shows a close match to the ordinary spectrum

from the single crystal.

Optical modeling results in best fits without significant parameter
correlation. Figure 54 shows the initial spectrum for a
CuIn~0.69~Ga~0.31~Se~2~ film obtained from the point-by-point fit. The
90% confidence limit values were less than 1% of the ε~2~ values so the
initial spectrum gave a good representation of the bulk optical
constants. The spectrum is free from the substrate effect and hence
there are no interference fringes. Also shown are the different
oscillators used to fit the initial spectrum. In order to fit the
initial spectra accurately, more oscillators (labeled A, B, and C) than
the number of the CP's in the measurement energy range were used. This
was necessitated by residual absorption due to electronic transitions at
higher energies, for example E~2~ (A), and also to account for possible
indirect transitions [^66]. The oscillators used in this model provided
a good fit for both ε~2~ and its derivatives. This approach of
generating the initial spectrum and subsequently fitting it with
oscillators helps to resolve minor details of the ε~2~ spectra that
cannot be resolved by directly fitting the ellipsometry parameters.
Table 38 lists the oscillator energies corresponding to the electronic
transitions in the measurement energy range along with the fit index
described in terms of MSE for all the films. The MSE values are close to
unity for low Ga films, indicating the high quality of the fits. MSE
values are slightly higher with higher Ga concentrations due to the
increase in the number interference fringes associated with the band gap
shift. If we restrict the fit to the absorbing region, the MSE value
would be close to unity. Table 37 also shows the comparison of the
surface roughness values obtained from the ellipsometry analysis and RMS
roughness from AFM measurements. The measurements show the same trend in
the roughness variation, although the values are slightly different.
This difference may be due to differences in the measurement techniques
or the sample region considered for each measurement. Figure 40(a) shows
the measured results and best fit for the ellipsometric parameters Ψ and
∆ measured on the CuIn~0.69~Ga~0.31~Se~2~ film at 70° incident angle.
The fit was performed at three incident angles, but the data for 65° and
75° incident angles are not shown here for clarity. The x-axis is shown
in terms of wavelength to highlight the fit in the transparent region
where the interference fringes dominate the spectrum. Figure 55(b) shows
the measurement and fit of the depolarization for the same film at 70^o^
incident angle. The depolarization spectrum shows interference fringes
similar to the Ψ and ∆ spectra, which are due to thickness
non-uniformity in the film. The depolarization spectrum is modeled with
3.5% non-uniformity in the CuIn~0.69~Ga~0.31~Se~2~ film thickness.
Experimental data in the transparent region were analyzed with extra
caution to separate the critical points E~0~(A), E~0~(B) and E~0~(C)
from interference fringes. Further detailed descriptions of the fit
strategies are beyond the scope of this report.

![](media/image65.wmf){width="3.4444444444444446in"
height="2.9722222222222223in"}

[]{#__RefHeading___Toc469989185 .anchor}Figure 54. Initial spectrum
(dashed line) for a typical CuIn~1-x~Ga~x~Se~2~

sample obtained from the point-by-point fit. The spectrum is free from

substrate effects as is clear from the absence of interference fringes.
Also

shown are the different oscillators used to fit the initial spectrum.

![](media/image66.wmf){width="3.0006944444444446in"
height="2.285416666666667in"}![](media/image67.wmf){width="2.71875in"
height="2.216666666666667in"}

[]{#__RefHeading___Toc469989186 .anchor}Figure 55. Measurement and fit
of the ellipsometry parameters

∆ and Ψ (a),and depolarization (b) for the CuIn~0.69~Ga~0.31~Se~2~ film

at 70^o^ incident angle. The depolarization is modeled with 3.5%
thickness

non-uniformity in the CuIn~0.69~Ga~0.31~Se~2~ film.

[]{#__RefHeading___Toc469989115 .anchor}Table 21. Oscillator energies
corresponds to the electronic transitions with varying Ga/(Ga+In) ratio.

+-----------+-------+-------+-------+-------+--------+-------+-------+
| G         | E~    | E~    | E~    | E~    | E (XΓ) | E~    | MSE   |
| a/(In+Ga) | 0~(A) | 0~(B) | 0~(C) | 1~(A) |        | 1~(B) |       |
|           |       |       |       |       | (eV)   |       |       |
|           | (eV)  | (eV)  | (eV)  | (eV)  |        | (eV)  |       |
+-----------+-------+-------+-------+-------+--------+-------+-------+
| 0.00      |       | 1.023 | 1.247 | 2.917 | 3.187  | 3.662 | 1.0   |
+-----------+-------+-------+-------+-------+--------+-------+-------+
| 0.31      | 1.168 | 1.208 | 1.449 | 2.965 | 3.302  | 3.746 | 1.7   |
+-----------+-------+-------+-------+-------+--------+-------+-------+
| 0.45      | 1.280 | 1.351 | 1.599 | 3.019 | 3.345  | 3.809 | 1.6   |
+-----------+-------+-------+-------+-------+--------+-------+-------+
| 0.66      | 1.366 | 1.457 | 1.641 | 3.066 | 3.379  | 3.917 | 2.6   |
+-----------+-------+-------+-------+-------+--------+-------+-------+
| 1.00      | 1.689 | 1.771 | 1.976 | 3.168 | 3.528  | 4.057 | 3.0   |
+-----------+-------+-------+-------+-------+--------+-------+-------+

Figure 56 shows the second derivatives of the ε~2~ spectrum for the
CuInSe~2~ film and the ordinary component of the CuInSe~2~ single
crystal (ε~2~) spectrum generated from *n* and *k* values in reference
45. The CP energies in the ε~2~ second derivative spectrum of the
CuInSe~2~ film match those of the CuInSe~2~ crystal. E~0~(A) and E~0~(B)
are not resolved in either spectra so the first transitions are marked
as E~0~(A,B). At the same time, the spin orbit splitting is well
resolved and the value ∆~so~ is 0.23eV consistent with the value
reported for a bulk CuInSe~2~ crystal [^67]. Higher energy interband
transitions corresponding to the electronic transitions E~1~(A) and
E~1~(B) at position N in the BZ and E(∆X) at position T in the BZ are
also visible and are shown in Figure 56. In general, the good
correlation between the film and crystal CP energies confirms the
integrity of the optical model used in this analysis.

Figures 57(a) and (b) show the optical constants *n* and *k* for
CuIn~1‑x~Ga~x~Se~2~ films with

0 ≤ x ≤ 1. Tabulated *n* and *k* values are listed in Table 22. The *n*
and *k* values are provided in 50 meV intervals over the range from
0.75eV to 2eV to resolve the spin orbit splitting, and in 100 meV
intervals from 2eV to 4.6eV. In general, ellipsometry measurements are
less sensitive to *k* values below the absorption band and published
single crystal optical data is not readily available for comparison in
this range. All electronic transitions shift to higher energy with
increasing Ga content. The spin orbit splitting is clearly visible in
the *n* and *k* spectra for the entire Ga range in the films. The
crystal field splitting is not distinguishable in the optical function
spectra but can be clearly seen in the second derivative of the ε~2~
spectra, which is shown in Figure 58. CP energies obtained from the
minima in the second derivative of the ε~2~ spectra are listed in Table
23. In general, the CP energies thus obtained are consistent with the
oscillator energies shown in Table 21. The crystal field splitting is
not resolved in the CuInSe~2~ films but becomes better resolved with
increasing Ga concentrations. This can be explained by the fact that the
distortion from cubic symmetry for CuInSe~2~ is small compared to
CuGaSe~2~ so the crystal field is not significant enough to resolve the
peaks at room temperature. High-resolution electro-reflectance
measurements on CuInSe~2~ gave ∆~so~=0.233eV and ∆~cf~ =0.006eV at 77K
\[Error: Reference source not found\]. Figure 59 shows the variation in
E~0~(A), E~0~(B) and E~0~(C) with Ga concentration and can be fit with a
2^nd^ degree polynomial as shown in equations 10-12. E~0~(A), for
CuInSe~2~ was determined using the electro-reflectance result for ∆~cf~
from the literature.

$E_{0}(A) = 1\text{.}\text{035} + 0\text{.}\text{389}x + 0\text{.}\text{264}x^{2}$
(10)

$E_{0}(B) = 1\text{.}\text{037} + 0\text{.}\text{626}x + 0\text{.}\text{104}x^{2}$
(11)

$E_{0}(C) = 1\text{.}\text{276} + 0\text{.}\text{644}x + 0\text{.}\text{057}x^{2}$
(12)

Since we assigned E~0~(A) to the lowest energy level according to
convention, there is no cross over between the energy levels E~0~(A) and
E~0~(B) as one would expect from the cross over between the valance
bands $\Gamma_{7v}^{4}$and $\Gamma_{6v}^{5}$ due to changes in the spin
orbit splitting. As Equation 10 shows, the optical bowing parameter for
the fundamental band gap is 0.264, which is slightly higher than the
bowing parameter of 0.21 based on a theoretical calculation for
stoichiometric material.^3^ However, the results based on experiments
were largely scattered 42 and varied from 0.11 [^68] to 0.24 [^69] for
evaporated films. This large discrepancy in the bowing parameter could
be partly due to the Cu stoichiometric deviation \[Error: Reference
source not found\].

![](media/image68.wmf){width="3.4027777777777777in"
height="2.513888888888889in"}

[]{#__RefHeading___Toc469989187 .anchor}Figure 56. Second derivatives of
the ε~2~ spectrum for a CuInSe~2~ film

(solid line) and the ordinary ε~2~ spectrum of a CuInSe~2~ single
crystal (dashed line). Critical points correspond to minima in the
spectra and are marked in the figure.

![](media/image69.wmf){width="2.926388888888889in"
height="2.214583333333333in"}
![](media/image70.wmf){width="2.9055555555555554in"
height="2.214583333333333in"}

[]{#__RefHeading___Toc469989188 .anchor}Figure 57. Index of refraction
(a) and extinction coefficient

\(b\) spectra for CuIn~1-x~Ga~x~Se~2~ films with varying Ga content.

The spin orbit splitting can be clearly seen.

![](media/image71.wmf){width="3.4027777777777777in"
height="3.3055555555555554in"}

[]{#__RefHeading___Toc469989189 .anchor}Figure 58. Second derivative of
the ε~2~ spectra for CuIn~1-x~Ga~x~Se~2~ films

with varying Ga content. All electronic transitions shift towards

higher energy with increasing Ga content.

![](media/image72.wmf){width="3.4027777777777777in"
height="2.6805555555555554in"}

[]{#__RefHeading___Toc469989190 .anchor}Figure 59. Variation in the
fundamental absorption transitions

E~0~(A), E~0~(B) and E~0~(C) of CuIn~1-x~Ga~x~Se~2~ films with x.

[]{#__RefHeading___Toc469989116 .anchor}Table 22. Optical constants *n*
and *k* for Cu(InGa)Se~2~ with varying x.

<table>
<colgroup>
<col style="width: 8%" />
<col style="width: 6%" />
<col style="width: 6%" />
<col style="width: 10%" />
<col style="width: 10%" />
<col style="width: 10%" />
<col style="width: 10%" />
<col style="width: 10%" />
<col style="width: 10%" />
<col style="width: 8%" />
<col style="width: 8%" />
</colgroup>
<thead>
<tr class="header">
<th rowspan="2"><p>E</p>
<p>(eV)</p></th>
<th colspan="2">x = 0</th>
<th colspan="2">x = 0.31</th>
<th colspan="2">x = 0.45</th>
<th colspan="2">x = 0.66</th>
<th colspan="2">x = 1.0</th>
</tr>
<tr class="odd">
<th><em>n</em></th>
<th><em>k</em></th>
<th><em>n</em></th>
<th><em>k</em></th>
<th><em>n</em></th>
<th><em>k</em></th>
<th><em>n</em></th>
<th><em>k</em></th>
<th><em>n</em></th>
<th><em>k</em></th>
</tr>
</thead>
<tbody>
<tr class="odd">
<td>4.60</td>
<td>2.81</td>
<td>1.26</td>
<td>2.73</td>
<td>1.33</td>
<td>2.80</td>
<td>1.30</td>
<td>2.78</td>
<td>1.32</td>
<td>2.67</td>
<td>1.30</td>
</tr>
<tr class="even">
<td>4.50</td>
<td>2.78</td>
<td>1.16</td>
<td>2.72</td>
<td>1.24</td>
<td>2.78</td>
<td>1.22</td>
<td>2.75</td>
<td>1.26</td>
<td>2.66</td>
<td>1.28</td>
</tr>
<tr class="odd">
<td>4.40</td>
<td>2.74</td>
<td>1.08</td>
<td>2.70</td>
<td>1.17</td>
<td>2.74</td>
<td>1.16</td>
<td>2.71</td>
<td>1.20</td>
<td>2.65</td>
<td>1.27</td>
</tr>
<tr class="even">
<td>4.30</td>
<td>2.69</td>
<td>1.02</td>
<td>2.67</td>
<td>1.12</td>
<td>2.71</td>
<td>1.12</td>
<td>2.68</td>
<td>1.18</td>
<td>2.66</td>
<td>1.27</td>
</tr>
<tr class="odd">
<td>4.20</td>
<td>2.64</td>
<td>0.99</td>
<td>2.63</td>
<td>1.09</td>
<td>2.68</td>
<td>1.10</td>
<td>2.66</td>
<td>1.18</td>
<td>2.68</td>
<td>1.27</td>
</tr>
<tr class="even">
<td>4.10</td>
<td>2.60</td>
<td>0.99</td>
<td>2.61</td>
<td>1.08</td>
<td>2.66</td>
<td>1.09</td>
<td>2.66</td>
<td>1.18</td>
<td>2.73</td>
<td>1.26</td>
</tr>
<tr class="odd">
<td>4.00</td>
<td>2.56</td>
<td>0.99</td>
<td>2.59</td>
<td>1.08</td>
<td>2.65</td>
<td>1.09</td>
<td>2.68</td>
<td>1.18</td>
<td>2.78</td>
<td>1.20</td>
</tr>
<tr class="even">
<td>3.90</td>
<td>2.54</td>
<td>1.02</td>
<td>2.59</td>
<td>1.09</td>
<td>2.66</td>
<td>1.10</td>
<td>2.72</td>
<td>1.17</td>
<td>2.77</td>
<td>1.13</td>
</tr>
<tr class="odd">
<td>3.80</td>
<td>2.54</td>
<td>1.05</td>
<td>2.61</td>
<td>1.10</td>
<td>2.69</td>
<td>1.09</td>
<td>2.74</td>
<td>1.12</td>
<td>2.74</td>
<td>1.11</td>
</tr>
<tr class="even">
<td>3.70</td>
<td>2.58</td>
<td>1.08</td>
<td>2.64</td>
<td>1.08</td>
<td>2.70</td>
<td>1.06</td>
<td>2.72</td>
<td>1.09</td>
<td>2.71</td>
<td>1.12</td>
</tr>
<tr class="odd">
<td>3.60</td>
<td>2.64</td>
<td>1.06</td>
<td>2.65</td>
<td>1.05</td>
<td>2.70</td>
<td>1.03</td>
<td>2.70</td>
<td>1.09</td>
<td>2.71</td>
<td>1.16</td>
</tr>
<tr class="even">
<td>3.50</td>
<td>2.64</td>
<td>1.03</td>
<td>2.64</td>
<td>1.04</td>
<td>2.68</td>
<td>1.04</td>
<td>2.69</td>
<td>1.13</td>
<td>2.77</td>
<td>1.22</td>
</tr>
<tr class="odd">
<td>3.40</td>
<td>2.63</td>
<td>1.02</td>
<td>2.63</td>
<td>1.05</td>
<td>2.69</td>
<td>1.06</td>
<td>2.73</td>
<td>1.18</td>
<td>2.85</td>
<td>1.27</td>
</tr>
<tr class="even">
<td>3.30</td>
<td>2.63</td>
<td>1.06</td>
<td>2.65</td>
<td>1.10</td>
<td>2.73</td>
<td>1.10</td>
<td>2.81</td>
<td>1.22</td>
<td>3.00</td>
<td>1.29</td>
</tr>
<tr class="odd">
<td>3.20</td>
<td>2.67</td>
<td>1.11</td>
<td>2.70</td>
<td>1.13</td>
<td>2.80</td>
<td>1.13</td>
<td>2.94</td>
<td>1.23</td>
<td>3.18</td>
<td>1.20</td>
</tr>
<tr class="even">
<td>3.10</td>
<td>2.74</td>
<td>1.15</td>
<td>2.79</td>
<td>1.16</td>
<td>2.91</td>
<td>1.13</td>
<td>3.10</td>
<td>1.17</td>
<td>3.31</td>
<td>1.02</td>
</tr>
<tr class="odd">
<td>3.00</td>
<td>2.86</td>
<td>1.18</td>
<td>2.92</td>
<td>1.14</td>
<td>3.04</td>
<td>1.06</td>
<td>3.21</td>
<td>1.01</td>
<td>3.30</td>
<td>0.84</td>
</tr>
<tr class="even">
<td>2.90</td>
<td>3.03</td>
<td>1.12</td>
<td>3.05</td>
<td>1.04</td>
<td>3.12</td>
<td>0.94</td>
<td>3.22</td>
<td>0.85</td>
<td>3.27</td>
<td>0.71</td>
</tr>
<tr class="odd">
<td>2.80</td>
<td>3.14</td>
<td>0.96</td>
<td>3.10</td>
<td>0.91</td>
<td>3.14</td>
<td>0.80</td>
<td>3.19</td>
<td>0.73</td>
<td>3.22</td>
<td>0.62</td>
</tr>
<tr class="even">
<td>2.70</td>
<td>3.14</td>
<td>0.80</td>
<td>3.10</td>
<td>0.79</td>
<td>3.12</td>
<td>0.70</td>
<td>3.15</td>
<td>0.64</td>
<td>3.18</td>
<td>0.55</td>
</tr>
<tr class="odd">
<td>2.60</td>
<td>3.11</td>
<td>0.71</td>
<td>3.07</td>
<td>0.71</td>
<td>3.09</td>
<td>0.62</td>
<td>3.11</td>
<td>0.58</td>
<td>3.14</td>
<td>0.50</td>
</tr>
<tr class="even">
<td>2.50</td>
<td>3.08</td>
<td>0.64</td>
<td>3.05</td>
<td>0.64</td>
<td>3.05</td>
<td>0.57</td>
<td>3.07</td>
<td>0.53</td>
<td>3.11</td>
<td>0.46</td>
</tr>
<tr class="odd">
<td>2.40</td>
<td>3.05</td>
<td>0.58</td>
<td>3.03</td>
<td>0.59</td>
<td>3.03</td>
<td>0.52</td>
<td>3.04</td>
<td>0.50</td>
<td>3.08</td>
<td>0.42</td>
</tr>
<tr class="even">
<td>2.30</td>
<td>3.02</td>
<td>0.54</td>
<td>3.01</td>
<td>0.55</td>
<td>3.00</td>
<td>0.49</td>
<td>3.02</td>
<td>0.47</td>
<td>3.06</td>
<td>0.39</td>
</tr>
<tr class="odd">
<td>2.20</td>
<td>3.00</td>
<td>0.51</td>
<td>2.99</td>
<td>0.52</td>
<td>2.98</td>
<td>0.46</td>
<td>3.01</td>
<td>0.44</td>
<td>3.05</td>
<td>0.36</td>
</tr>
<tr class="even">
<td>2.10</td>
<td>2.99</td>
<td>0.48</td>
<td>2.98</td>
<td>0.49</td>
<td>2.97</td>
<td>0.43</td>
<td>2.99</td>
<td>0.42</td>
<td>3.04</td>
<td>0.34</td>
</tr>
<tr class="odd">
<td>2.00</td>
<td>2.98</td>
<td>0.46</td>
<td>2.97</td>
<td>0.47</td>
<td>2.96</td>
<td>0.41</td>
<td>2.98</td>
<td>0.39</td>
<td>3.04</td>
<td>0.31</td>
</tr>
<tr class="even">
<td>1.95</td>
<td>2.97</td>
<td>0.45</td>
<td>2.97</td>
<td>0.45</td>
<td>2.96</td>
<td>0.40</td>
<td>2.98</td>
<td>0.38</td>
<td>3.05</td>
<td>0.28</td>
</tr>
<tr class="odd">
<td>1.90</td>
<td>2.97</td>
<td>0.43</td>
<td>2.97</td>
<td>0.44</td>
<td>2.96</td>
<td>0.39</td>
<td>2.98</td>
<td>0.37</td>
<td>3.04</td>
<td>0.24</td>
</tr>
<tr class="even">
<td>1.85</td>
<td>2.96</td>
<td>0.42</td>
<td>2.96</td>
<td>0.43</td>
<td>2.95</td>
<td>0.38</td>
<td>2.99</td>
<td>0.36</td>
<td>3.02</td>
<td>0.23</td>
</tr>
<tr class="odd">
<td>1.80</td>
<td>2.96</td>
<td>0.41</td>
<td>2.97</td>
<td>0.42</td>
<td>2.96</td>
<td>0.37</td>
<td>2.99</td>
<td>0.35</td>
<td>3.03</td>
<td>0.22</td>
</tr>
<tr class="even">
<td>1.75</td>
<td>2.96</td>
<td>0.40</td>
<td>2.97</td>
<td>0.40</td>
<td>2.96</td>
<td>0.35</td>
<td>3.00</td>
<td>0.33</td>
<td>3.04</td>
<td>0.18</td>
</tr>
<tr class="odd">
<td>1.70</td>
<td>2.96</td>
<td>0.39</td>
<td>2.97</td>
<td>0.39</td>
<td>2.96</td>
<td>0.34</td>
<td>3.01</td>
<td>0.31</td>
<td>3.03</td>
<td>0.14</td>
</tr>
<tr class="even">
<td>1.65</td>
<td>2.96</td>
<td>0.38</td>
<td>2.97</td>
<td>0.38</td>
<td>2.97</td>
<td>0.33</td>
<td>3.01</td>
<td>0.28</td>
<td>2.99</td>
<td>0.08</td>
</tr>
<tr class="odd">
<td>1.60</td>
<td>2.96</td>
<td>0.37</td>
<td>2.98</td>
<td>0.37</td>
<td>2.97</td>
<td>0.31</td>
<td>3.00</td>
<td>0.26</td>
<td>2.94</td>
<td>0.07</td>
</tr>
<tr class="even">
<td>1.55</td>
<td>2.97</td>
<td>0.36</td>
<td>2.98</td>
<td>0.36</td>
<td>2.98</td>
<td>0.29</td>
<td>3.00</td>
<td>0.24</td>
<td>2.92</td>
<td>0.07</td>
</tr>
<tr class="odd">
<td>1.50</td>
<td>2.97</td>
<td>0.35</td>
<td>3.00</td>
<td>0.34</td>
<td>2.99</td>
<td>0.27</td>
<td>3.02</td>
<td>0.22</td>
<td>2.89</td>
<td>0.06</td>
</tr>
<tr class="even">
<td>1.45</td>
<td>2.98</td>
<td>0.34</td>
<td>3.01</td>
<td>0.32</td>
<td>2.98</td>
<td>0.24</td>
<td>3.04</td>
<td>0.16</td>
<td>2.87</td>
<td>0.06</td>
</tr>
<tr class="odd">
<td>1.40</td>
<td>2.99</td>
<td>0.33</td>
<td>3.02</td>
<td>0.28</td>
<td>2.98</td>
<td>0.22</td>
<td>2.99</td>
<td>0.11</td>
<td>2.86</td>
<td>0.06</td>
</tr>
<tr class="even">
<td>1.35</td>
<td>3.00</td>
<td>0.32</td>
<td>3.01</td>
<td>0.26</td>
<td>3.00</td>
<td>0.20</td>
<td>2.99</td>
<td>0.09</td>
<td>2.84</td>
<td>0.05</td>
</tr>
<tr class="odd">
<td>1.30</td>
<td>3.01</td>
<td>0.30</td>
<td>3.02</td>
<td>0.24</td>
<td>3.01</td>
<td>0.15</td>
<td>2.94</td>
<td>0.04</td>
<td>2.83</td>
<td>0.05</td>
</tr>
<tr class="even">
<td>1.25</td>
<td>3.03</td>
<td>0.27</td>
<td>3.03</td>
<td>0.21</td>
<td>2.98</td>
<td>0.09</td>
<td>2.90</td>
<td>0.04</td>
<td>2.82</td>
<td>0.05</td>
</tr>
<tr class="odd">
<td>1.20</td>
<td>3.02</td>
<td>0.24</td>
<td>3.05</td>
<td>0.15</td>
<td>2.94</td>
<td>0.06</td>
<td>2.87</td>
<td>0.04</td>
<td>2.80</td>
<td>0.04</td>
</tr>
<tr class="even">
<td>1.15</td>
<td>3.02</td>
<td>0.22</td>
<td>3.01</td>
<td>0.08</td>
<td>2.90</td>
<td>0.05</td>
<td>2.85</td>
<td>0.04</td>
<td>2.79</td>
<td>0.04</td>
</tr>
<tr class="odd">
<td>1.10</td>
<td>3.03</td>
<td>0.21</td>
<td>2.96</td>
<td>0.06</td>
<td>2.87</td>
<td>0.05</td>
<td>2.83</td>
<td>0.03</td>
<td>2.78</td>
<td>0.04</td>
</tr>
<tr class="even">
<td>1.05</td>
<td>3.07</td>
<td>0.18</td>
<td>2.92</td>
<td>0.05</td>
<td>2.85</td>
<td>0.04</td>
<td>2.82</td>
<td>0.03</td>
<td>2.77</td>
<td>0.04</td>
</tr>
<tr class="odd">
<td>1.00</td>
<td>3.05</td>
<td>0.05</td>
<td>2.90</td>
<td>0.05</td>
<td>2.84</td>
<td>0.04</td>
<td>2.81</td>
<td>0.03</td>
<td>2.77</td>
<td>0.03</td>
</tr>
<tr class="even">
<td>0.95</td>
<td>2.98</td>
<td>0.04</td>
<td>2.88</td>
<td>0.05</td>
<td>2.82</td>
<td>0.04</td>
<td>2.79</td>
<td>0.03</td>
<td>2.76</td>
<td>0.03</td>
</tr>
<tr class="odd">
<td>0.90</td>
<td>2.95</td>
<td>0.04</td>
<td>2.86</td>
<td>0.04</td>
<td>2.81</td>
<td>0.03</td>
<td>2.78</td>
<td>0.03</td>
<td>2.75</td>
<td>0.03</td>
</tr>
<tr class="even">
<td>0.85</td>
<td>2.92</td>
<td>0.04</td>
<td>2.85</td>
<td>0.04</td>
<td>2.80</td>
<td>0.03</td>
<td>2.77</td>
<td>0.02</td>
<td>2.74</td>
<td>0.03</td>
</tr>
<tr class="odd">
<td>0.80</td>
<td>2.90</td>
<td>0.04</td>
<td>2.84</td>
<td>0.04</td>
<td>2.79</td>
<td>0.03</td>
<td>2.76</td>
<td>0.02</td>
<td>2.74</td>
<td>0.03</td>
</tr>
<tr class="even">
<td>0.75</td>
<td>2.89</td>
<td>0.03</td>
<td>2.83</td>
<td>0.03</td>
<td>2.78</td>
<td>0.03</td>
<td>2.76</td>
<td>0.02</td>
<td>2.73</td>
<td>0.02</td>
</tr>
</tbody>
</table>

The spin orbit splitting and crystal field splitting were determined
using Equations 6 and 7 and are shown in Figure 60. The ordering of the
valance band $\Gamma_{7v}^{4}$and $\Gamma_{6v}^{5}$ was determined from
the crystal data for similar compositions reported in the
literatureError: Reference source not found. The error bars shown are
the SE measurement resolution due to the energy step used for this
study. The ∆~so~ shows an upward bowing and the values are consistent
with other results reported in the literature [^70]. Small differences
in the data could be due to the stoichiometry deviations in the samples.
The negative values for ∆~cf~ are in accordance with the polarization
selection rule and negative crystal field observed in the
CuIn~1-x~Ga~x~Se~2~ ternary chalcopyrites. [^71], [^72]~.~ The crystal
field splitting shows a non-linear dependence with the Ga concentration.
Vegard's law suggests a linear variation in the c/a ratio with Ga
content, which results in a linear variation in the crystal field
splitting based on the compression rule as described in the Equation 4.
However, the nonlinear behavior observed in the data may be caused by
the fact that for Cu-based ternary chalcopyrites the crystal field
splitting is not purely dependent on the tetragonal distortion41 Band
structure calculations show that the optical band gap is also influenced
by the anion displacement \[Error: Reference source not found\]. The
accuracy of crystal field splitting data is limited by the sample
nature; slightly textured polycrystalline films. For more accurate
values, measurements on single crystal or epitaxial films are
recommended.

[]{#__RefHeading___Toc469989117 .anchor}Table 23. Variation in the
electronic transition energies with Ga/(Ga+In) ratio obtained from
Figure 59.

+----------+---------+---------+---------+---------+---------+---------+
| Ga       | E~0~(A) | E~0~(B) | E~0~(C) | E~1~(A) | E(XΓ)   | E~1~(B) |
| /(In+Ga) |         |         |         |         |         |         |
|          | (eV)    | (eV)    | (eV)    | (eV)    | (eV)    | (eV)    |
+----------+---------+---------+---------+---------+---------+---------+
| 0.00     |         | 1.03    | 1.27    | 2.91    | 3.19    | 3.66    |
+----------+---------+---------+---------+---------+---------+---------+
| 0.31     | 1.18    | 1.23    | 1.48    | 2.96    | 3.31    | 3.74    |
+----------+---------+---------+---------+---------+---------+---------+
| 0.45     | 1.27    | 1.35    | 1.58    | 3.02    | 3.34    | 3.84    |
+----------+---------+---------+---------+---------+---------+---------+
| 0.66     | 1.40    | 1.49    | 1.72    | 3.07    | 3.37    | 3.91    |
+----------+---------+---------+---------+---------+---------+---------+
| 1.00     | 1.69    | 1.77    | 1.98    | 3.17    | 3.52    | 4.06    |
+----------+---------+---------+---------+---------+---------+---------+

Higher interband transitions are also visible in the ε~2~ spectra.
Figure 61 shows the effect of Ga variation on transitions E~1~(A), E(X)
and E~1~(B). E(X) corresponds to the transition at the BZ center from
~5v~^(2)^ to ~1c~. E~1~(A) and E~1~(B) correspond to transitions at
the edge point N in the BZ. E(∆X), corresponding to an electronic
transition at T in the BZ, is not clearly visible with all compositions.

### 3.3.7 Cu(InGa)Se~2~ Ellipsometry: Conclusions

The surface roughness of polycrystalline CuIn~1‑x~Ga~x~Se~2~ films
deposited by four-source co-evaporation made them unsuitable for SE
analysis. An innovative method was developed to circumvent this problem
by performing the SE measurement on the reverse side of the films after
they were peeled off from the substrate. The films peeled cleanly from
the substrate, probably due to the presence of MoSe~2~, and smooth
crack-free CuIn~1‑x~Ga~x~Se~2~ films were obtained. The presence of an
optically thick Mo film, deposited prior to the peeling, made the
optical modeling easier. Also RMS roughness from AFM measurements
performed after Mo deposition provided an estimation of the interface
layer thickness.

A detailed description of the SE measurement and multilayer modeling
strategy generic to ternary chalcopyrite semiconductors was presented.
The optical functions hence obtained are compared with the data reported
in the literature from single /polycrystalline ingots. Optical constants
measured on polycrystalline CuIn~1‑x~Ga~x~Se~2~ films were found to be
similar to the ordinary component of the single crystal optical
constants. A comparison of the ε~2~ spectra for a CuInSe~2~ film and
single crystal shows that the polycrystalline nature of the film did not
affect the optical constants. This indicates that the change in the path
length as a result of grain boundary reflection is not significant for
these films. In other words, the evaporated CuIn~1‑x~Ga~x~Se~2~ films
are compact and large grained. SE measurements in the anisotropy mode
show that the deviation from the random orientation did not cause any
measurable anisotropy in the film.

The effect of Ga concentration on different optical transitions in the
dielectric spectra was shown in detail. As the Ga concentration
increases, the electronic transitions shift to higher energies
consistent with the theoretical and experimental results reported in the
literature. CuIn~1‑x~Ga~x~Se~2~ thin films show a triply degenerate
fundamental transition consistent with the chalcopyrite band structure.
The degeneracy was removed by crystal field splitting and spin-orbit
interaction. CuIn~1‑x~Ga~x~Se~2~ films show negative crystal field
splitting. For CuInSe~2~ the crystal field splitting is negligibly small
and is not measurable with room temperature SE measurements. As the Ga
concentration increases, the crystal field splitting increased similar
to the results reported for polycrystalline ingots in the literature.
Non-linear behavior was observed, in contradiction to the linear
behavior expected from Vegard\'s law and the compression rule. The
reason for this non-linear behavior could be the anion displacement as a
result of the two captions in the CuIn~1‑x~Ga~x~Se~2~ lattice. The
bowing parameter for the band gap is slightly higher than the values
reported in the literature based on theoretical calculation. Second
order polynomial equations are presented for the Ga dependence for the
fundamental electronic transitions. Tabulated values of optical
constants.

### 3.3.8 Team Activities and Collaboration

### 

Bill Shafarman attended the National CIS R&D Team meeting on January
30-31, 2003 where he participated in work by the Absorber Sub-team and
presented a review titled "Status of Wide Bandgap CuInSe~2~-based Solar
Cells". Additional team-related activity and collaboration included:

1.  Cu(InGa)Se~2~ samples deposited at either 400°C or 550°C using
    either a uniform or bi-layer (Cu-rich at start) deposition process
    were provided to the University of Illinois for TEM
    characterization. A set of well characterized Cu(InGa)Se~2~ films
    were also provided to UI for EDS standards.

2.  IEC fabricated devices on a set of 12 Cu(InGa)S~2~ films from
    Florida Solar Energy Center and completed J-V and QE measurements to
    assist FSEC in evaluating it's absorber layer and device fabrication
    processes.

3.  Optical constants measured by spectroscopic ellipsometry as a
    function of energy of Cu(InGa)Se~2~ with different relative Ga
    content were sent to several team members for input into their
    modeling efforts.

4.  CdS and ZnO/ITO depositions were completed on several sets of
    samples, in a round-robin experiment with GSE and NREL, to use for
    characterization of GSE's device processing and calibration of their
    emitter and TCO layer thicknesses.

5.  Collaboration with the University of Oregon on device
    characterization continued with focus on Cu(InAl)Se~2~ devices with
    different bandgaps for comparison to previously characterized
    Cu(InGa)Se~2~ devices.

![](media/image73.wmf){width="3.763888888888889in"
height="2.6805555555555554in"}

[]{#__RefHeading___Toc469989191 .anchor}Figure 60. Variation in the
∆~cf~ and ∆~so~ of CuIn~1-x~Ga~x~Se~2~ films with x.

The non-linear behavior of the crystal field splitting could be due

to the effect of anion displacement.

![](media/image74.wmf){width="3.4027777777777777in"
height="2.513888888888889in"}

[]{#__RefHeading___Toc469989192 .anchor}Figure 61. Variation in the
higher energy transitions of CuIn~1-x~Ga~x~Se~2~ films with x.

# 

# **4 SI BASED SOLAR CELLS**

## 4.1 Fabrication of Large Grain Si Films

The grain size enhancement of HWCVD Si films deposited on Al layers has
been investigated in two temperature regimes. The eutectic temperature,
575°C, determines these regimes in the Si-Al system. With the formation
of a liquid phase above the eutectic temperature, a different mechanism
is expected to govern the grain enhancement process compared to that
expected below this temperature where the Si-Al mixture remains in the
solid state. The following sections describe the results obtained for
depositions of HWCVD Si films on Al layers at 430°C and 600°C.

### 4.1.1 Grain size enhancement of HWCVD Si films at 430°C.

This work was published by Ebil et. al. [^73]. Aluminum layers with
thickness varying from 0.01 to 1.0 μm were deposited onto Corning 7059
type glass substrates using electron-beam evaporation. Silicon layers
were subsequently deposited onto the Al coated samples in the HWCVD
system. Silane was used as the source gas with no hydrogen dilution.
Silicon layer depositions were carried out at a substrate temperature of
430°C, filament temperature of 1850°C and a pressure of 25m Torr. The
films were deposited at two different growth rates of 1.0 and 5.0 μm/hr,
which were attained by varying the silane flow rate from 4.5 to 22.5
sccm, respectively. Films with thickness values of 2 and 5 μm were
deposited with deposition times corresponding to 24 minutes and 5 hrs,
respectively.

Figure 62 shows a comparison of typical XRD patterns of films deposited
on glass and on an Al layers. All films deposited on Al had random
orientation and higher degree of crystallization compared to films
deposited on glass, as evidenced by the number of peaks appearing in the
pattern. This observation is corroborated by Raman spectroscopy
analysis, as shown in Figure 63. The crystalline Si Raman profile of
films deposited on glass had a typical tail extending toward lower
wavenumbers. This is indicative of a-Si and nanocrystalline Si phases
being present. In contrast, films deposited on Al, had profiles similar
to those obtained for c-Si. Table 24 summarizes the experimental
parameter variation and film properties measured. The average grain size
of the films was estimated from the broadening of the strongest Si XRD
peak and the crystalline fraction (X~c~) was calculated from the
integrated areas of the Raman peaks at 520 cm^-1^, the single-crystal
mode, at 504 cm^-1^, the nano-crystalline mode, and at 480 cm^-1^, the
amorphous mode, using the formula X~c~ = A~520~/(A~520~+A~504~+A~480~).
Since the Raman measurements sample approximately the top micron of the
film, it did not reflect the complete film structure for films with
thickness greater than this value. The crystalline fraction for the
films discussed here is, in general, the fraction of large grains, \>
100 nm, to nano-crystalline size grains.

![](media/image75.jpeg){width="6.002777777777778in"
height="4.001388888888889in"}

[]{#__RefHeading___Toc469989193 .anchor}Figure 62. XRD pattern of
poly-Si on glass, and poly-Si on Al coated glass.

For 2 μm thick films grown at 1 μm/hr, the average grain size and
crystalline fraction increases with decreasing Si/Al ratio, i.e., with
increasing Al layer thickness. For this particular set of films, the
maximum average grain size of 200 nm was achieved with a Si/Al ratio of
2. It is important to mention that this maximum grain size value
corresponds to the measurement limit of the XRD peak-broadening method
as imposed by the instrumental function of the diffractometer used.
Therefore, for films with such level of broadening, the average grain
size may be larger than 200 nm. For 2 μm thick films as the growth rate
is increased from 1 to 5 μm/hr, the average grain size and crystalline
fraction decreases at a constant Si/Al ratio. AFM analysis revealed that
some areas on the film surface are covered with small particles, 10-50
nm in diameter, as shown in Figure 64. It is speculated that these
features are individual grains, which do not extend throughout the bulk
of the Si film. However, these features were not found throughout the
film surface, which may indicate the presence of a large grain
microstructure. While AFM can only give details about the surface
morphology, XRD analysis gives information about the average grain size
of the bulk film. TEM analysis was carried out to investigate the
microstructure of a 2 μm thick sample deposited at 1.0 μm/hr growth rate
with a Si/Al ratio of 2. The cross-sectional TEM images of this sample
are shown in Figure 65. Several large grains and nano-crystalline
regions can be identified. These two crystalline regions can be found
either side by side on the Al layer, or as layers with one evolving from
the other. It is speculated that the presence of the nano-crystalline
phase indicates an incomplete grain enhancement process and that the
effect of lower growth rate for the given film thickness is to increase
the interaction time between the layers. Glancing-angle XRD analysis
showed traces of Al on the surface of

2 μm thick films. The bulk of the films also contain Al. Figure 66 is a
planar SEM image of one of this film after etching with 3:3:1
HF:HNO~3~:CH~3~COOH solution. Clearly, two microstructures are revealed.
The overlying layer corresponds to Al rich Si-Al alloy, while the
underlying structure is composed of pure poly-Si. The image appears to
indicate that the poly-Si does not form a continuous layer. This
conclusion is corroborated by optical microscopy and Raman spectroscopy
as shown in Figure 67. In this image of the back of the sample and
obtained through the glass substrate, the dark areas correspond to
poly-Si while the bright areas correspond to Al. It is evident that the
Al layer inversion has not reached completion, perhaps due to
insufficient interaction time between the Si and the Al layer. For 5.0
μm thick samples the average grain size varied between 40 and 70 nm, and
the crystalline fraction did not change significantly with the growth
rate. The decrease in grain size for Si films with 5 μm thickness may be
due to the existence of a limiting thickness beyond which the Al ceases
to have an effect on the growth process. If this hypothesis is correct,
the film evolves from a large grain to a nano-crystalline grain
microstructure as seen on films deposited on glass. All of the 5.0 μm
thick films deposited at 5.0 μm/hr growth rate had columnar features 4-5
μm in diameter. These features were uniformly distributed throughout the
sample. SEM images showed that some of these features were 3-4 μm high
from the surface. Although those features looked like individual large
grains, detailed SEM and AFM analysis revealed that they are actually
agglomerates of much smaller crystallites, 10-50 nm in size. At 1.0
μm/hr growth rate, the columnar features were absent and the film
surface appeared smoother in comparison.

![](media/image76.png){width="3.75in" height="3.8020833333333335in"}

[]{#__RefHeading___Toc469989194 .anchor}Figure 63. Comparison of Raman
spectra of c-Si, poly-Si on glass,

and poly-Si on Al coated glass.

[]{#__RefHeading___Toc469989118 .anchor}Table 24. Properties of films
deposited at 430°C.

+---+--------+----+-------+----------+--------+-----+-----+----------+
| > | Growth |    |       |          | Growth |     |     |          |
|   | rate 1 |    |       |          | rate 5 |     |     |          |
| 2 | μm/hr  |    |       |          | μm/hr  |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| μ |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| S |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| f |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| l |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| s |        |    |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | Thi    | S  | Avg.  | Cry      | Thi    | Si  | A   | Cry      |
|   | ckness | i/ |       | stalline | ckness | /Al | vg. | stalline |
|   |        | Al | grain | vol.     |        |     |     | vol.     |
|   | Al     |    |       | frac. at | Al     |     | gr  | frac. at |
|   | layer  |    | size  | 520      | layer  |     | ain | 520      |
|   |        |    | (nm)  | cm^-1^   |        |     |     | cm^-1^   |
|   | (μm)   |    |       |          | (μm)   |     | s   |          |
|   |        |    |       | (%)      |        |     | ize | (%)      |
|   |        |    |       |          |        |     | (   |          |
|   |        |    |       |          |        |     | nm) |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.01   | 2  | 60    | 85       | 0.01   | 200 | 40  | 65       |
|   |        | 00 |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.1    | 20 | 70    | 89       | 0.1    | 20  | 45  | 70       |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.5    | 4  | 100   | 96       | 0.5    | 4   | 70  | 71       |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 1      | 2  | 200   | 98       | 1      | 2   | 80  | 76       |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   |        |    |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
| > | Growth |    |       |          | Growth |     |     |          |
|   | rate 1 |    |       |          | rate 5 |     |     |          |
| 5 | μm/hr  |    |       |          | μm/hr  |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| μ |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| S |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| f |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| l |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| s |        |    |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | Thi    | S  | Avg.  | Cry      | Thi    | Si  | A   | Cry      |
|   | ckness | i/ |       | stalline | ckness | /Al | vg. | stalline |
|   |        | Al | grain | vol.     |        |     |     | vol.     |
|   | Al     |    |       | frac. at | Al     |     | gr  | frac. at |
|   | layer  |    | size  | 520      | layer  |     | ain | 520      |
|   |        |    | (nm)  | cm^-1^   |        |     |     | cm^-1^   |
|   | (μm)   |    |       |          | (μm)   |     | s   |          |
|   |        |    |       | (%)      |        |     | ize | (%)      |
|   |        |    |       |          |        |     | (   |          |
|   |        |    |       |          |        |     | nm) |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.01   | 5  | 40    | 77       | 0.01   | 500 | 40  | 71       |
|   |        | 00 |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.1    | 50 | 48    | 77       | 0.1    | 50  | 47  | 76       |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.5    | 10 | 54    | 79       | 0.5    | 10  | 66  | 73       |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 1      | 5  | 60    | 79       | 1      | 5   | 70  | 78       |
+---+--------+----+-------+----------+--------+-----+-----+----------+

N/A = Not Applicable N/M = Not Measured

![](media/image77.jpeg){width="3.567361111111111in"
height="2.9972222222222222in"}

[]{#__RefHeading___Toc469989195 .anchor}Figure 64. AFM image of the
surface of a 2 μm thick poly-Si film deposited Al.

![](media/image78.jpeg){width="5.195138888888889in" height="1.70625in"}

[]{#__RefHeading___Toc469989196 .anchor}Figure 65. XTEM image of 2 μm
thick poly-Si film on Al coated glass.

> ![](media/image79.jpeg){width="3.40625in"
> height="2.661111111111111in"}

[]{#__RefHeading___Toc469989197 .anchor}Figure 66. SEM planar image
after etching Si film deposited on Al.

![](media/image80.jpeg){width="3.602777777777778in"
height="2.7944444444444443in"}

[]{#__RefHeading___Toc469989198 .anchor}Figure 67. Optical micrograph of
back surface of Si film deposited on Al.

### 4.1.2 Grain size enhancement of HWCVD Si films at 600°C.

The deposition of these films was carried out as described in Section
4.1.1, with the exception of the substrate temperature. Film properties
as a function of deposition parameters are shown in Table 25. Compared
to films deposited at 430°C, films deposited at 600°C exhibits larger
average grain size as measured by XRD peak broadening. This contrast is
more pronounced as the growth rate and Si film thickness increases. Only
Si films with 5 μm thickness and deposited at 1.0 μm/hr showed random or
(220) orientation while all others had strong (111) orientation.
However, there is not enough data yet to correlate the orientation to
either growth rate or film thickness. All films deposited at 600°C were
non-homogeneous in terms of morphology and composition across their
surface, as shown in Figure 68. As a result, the crystalline fraction
reported in Table 25 is given as a range, with the larger crystalline
fraction corresponding to the areas with larger Al concentration in the
film. The non-homogeneity observed in these films could be the result of
poor wetting of the Al layer on the 7059 glass at 600°C or alternatively
due to segregation of the liquid phase forming at the eutectic
composition. On average, the crystalline volume fraction of films
deposited at 600°C is similar to those deposited at 430°C. However, as
shown in Figure 69, the Raman profiles of regions rich in Al indicate a
crystalline quality similar to that of c-Si for all films, regardless of
deposition conditions.

[]{#__RefHeading___Toc469989119 .anchor}Table 25. Properties of films
Deposited at 600°C.

+---+--------+----+-------+----------+--------+-----+-----+----------+
| > | Growth |    |       |          | Growth |     |     |          |
|   | rate 1 |    |       |          | rate 5 |     |     |          |
| 1 | μm/hr  |    |       |          | μm/hr  |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| μ |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| S |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| f |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| l |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| s |        |    |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | Thi    | S  | Avg.  | Cry      | Thi    | Si  | A   | Cry      |
|   | ckness | i/ |       | stalline | ckness | /Al | vg. | stalline |
|   |        | Al | grain | vol.     |        |     |     | vol.     |
|   | Al     |    |       | frac. at | Al     |     | gr  | frac. at |
|   | layer  |    | size  | 520      | layer  |     | ain | 520      |
|   |        |    | (nm)  | cm^-1^   |        |     |     | cm^-1^   |
|   | (μm)   |    |       |          | (μm)   |     | s   |          |
|   |        |    |       | (%)      |        |     | ize | (%)      |
|   |        |    |       |          |        |     | (   |          |
|   |        |    |       |          |        |     | nm) |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0      | N  | 40    | 46       | 0      | N/A | 30  | 30       |
|   |        | /A |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.1    | 10 | 130   | 70-98    | 0.1    | 10  | 90  | 70-95    |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.5    | 2  | 180   | 70-98    | 0.5    | 2   | 160 | 70-95    |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 1      | 1  | 160   | 70-98    | 1      | 1   | 190 | 70-95    |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   |        |    |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
| > | Growth |    |       |          | Growth |     |     |          |
|   | rate 1 |    |       |          | rate 5 |     |     |          |
| 5 | μm/hr  |    |       |          | μm/hr  |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| μ |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| S |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| > |        |    |       |          |        |     |     |          |
|   |        |    |       |          |        |     |     |          |
| f |        |    |       |          |        |     |     |          |
| i |        |    |       |          |        |     |     |          |
| l |        |    |       |          |        |     |     |          |
| m |        |    |       |          |        |     |     |          |
| s |        |    |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | Thi    | S  | Avg.  | Cry      | Thi    | Si  | A   | Cry      |
|   | ckness | i/ |       | stalline | ckness | /Al | vg. | stalline |
|   |        | Al | grain | vol.     |        |     |     | vol.     |
|   | Al     |    |       | frac. at | Al     |     | gr  | frac. at |
|   | layer  |    | size  | 520      | layer  |     | ain | 520      |
|   |        |    | (nm)  | cm^-1^   |        |     |     | cm^-1^   |
|   | (μm)   |    |       |          | (μm)   |     | s   |          |
|   |        |    |       | (%)      |        |     | ize | (%)      |
|   |        |    |       |          |        |     | (   |          |
|   |        |    |       |          |        |     | nm) |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0      | N  | 36    | 45       | 0      | N/A | 40  | 40       |
|   |        | /A |       |          |        |     |     |          |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.1    | 50 | 80    | 56-95    | 0.1    | 50  | N/M | 60-95    |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 0.5    | 10 | 130   | 56-95    | 0.5    | 10  | 110 | 60-95    |
+---+--------+----+-------+----------+--------+-----+-----+----------+
|   | 1      | 5  | 170   | 56-95    | 1      | 5   | 170 | 60-95    |
+---+--------+----+-------+----------+--------+-----+-----+----------+

N/A = Not Applicable

N/M = Not Measured

> ![](media/image81.jpeg){width="4.99375in"
> height="2.7729166666666667in"}

[]{#__RefHeading___Toc469989199 .anchor}Figure 68. Planar SEM/EDS of Si
film deposited on Al at 600°C.

![](media/image82.png){width="4.552083333333333in"
height="3.8020833333333335in"}

[]{#__RefHeading___Toc469989200 .anchor}Figure 69. Raman profiles of Al
rich regions in films deposited at 600°C.

## 4.2 Solar Cell Development 

### 4.2.1 Microcrystalline p-i-n solar cell by PECVD

An attempt was made to fabricate p-i-n solar cells with all PECVD µc-Si
layers. The device was deposited on glass/SnO~2~ substrates (LOF Tec 8
and Tec 15). Conditions were nearly the same for all three layers:
T=175°C (150°C for the p-layer), P=0.7 Torr (1.0 Torr for the p-layer),
and flow ratio H~2~/SiH~4~=180 sccm/2 sccm=90. The 2.0 µm i-layer was
deposited for 9 hours. No a-Si buffer layers or any other transitional
layers were incorporated. Device performance was very poor, with
V~oc~=0.2V and efficiency \<1%. Capacitance-voltage measurements
indicated the i-layer was 2 µm thick as planned. A strong frequency
dependency and unusual voltage bias dependence suggested a relatively
high defect density and low barrier height. Spectral response
measurements were made at 0 and --2V. The integrated photocurrent
increased from 4 to 8 mA/cm^2^ with reverse bias, confirming the poor
carrier collection, in the absence of an electric field. Figure 70 shows
the log QE vs energy for this device (GD41011) compared to a 10%
efficient a-Si device (GD4666), a HWCVD p-i-n device with a µc-Si
i-layer (HW193), and a 300 µm standard c-Si p-n cell. Both of the µc-Si
devices have substantially more absorption below the a-Si absorption
edge, confirming that their structure is between a-Si and c-Si.

![](media/image83.wmf){width="3.8444444444444446in"
height="3.517361111111111in"}

[]{#__RefHeading___Toc469989201 .anchor}Figure 70. Log QE vs energy for
a-Si, µc-Si and c-Si devices.

Data from HWCVD and PECVD µc-Si p-i-n cells is shown.

### 

### 4.2.2 Solar Cells with µc-Si HWCVD or PECVD i-layers

Thin film Si n-i-p solar cells were fabricated on (100) and (111)
n-doped c-Si substrates using PECVD and HWCVD microcrystalline i-layers.
No metal induced crystallization was utilized. The structure of the
devices was Ag-Pd-Ti/(100) or (111)n^+^-cSi/(PECVD or HWCVD)
i-μcSi/PECVD p-μcSi/ITO/Ni-Al. The n+ wafer substrates served as the
n-layer in the p-i-n cells. In addition, diagnostic n+-p+ devices were
fabricated by depositing PECVD p-μcSi emitter layers directly onto the
n^+^-c-Si substrates similar to the n+-p+ cell structure. Identical
n-i-p and n+-p+ devices were also fabricated which incorporated a 30 nm
a-Si buffer layer between the p-layer and i-layer or c-Si. Tables 26 and
27 summarize the deposition conditions used for the Si i-layers and the
a-Si buffer and p-layers.

The performance of all solar cell devices was quite poor. In general,
all cells with a HW or GD i-layer absorber had less than 1% efficiency
although many had well formed diode characteristics in light and dark.
The best devices were on pieces that consisted of only n+/p+ with no
µc-Si HW or GD absorbers. Table 28 shows analysis of the dark JV curves
on several devices. Only the Si layers are listed, all devices had same
wafer metallization and ITO/grid contacts. The 30 nm intrinsic buffer
layer is indicated as just a-Si. Intrinsic µc-Si absorbers indicated as
HW i or GD i. W is depletion width at 0V, N is carrier density from
1/C^2^ vs V. All devices had very low series resistance (\<0.5 Ω/cm^2^).
GD41015-21 (n+-p+ cell with no buffer) had A\~2, while the others had A
much greater than 2. All devices had very high J~0~ indicating high
recombination.

[]{#__RefHeading___Toc469989120 .anchor}Table 26. Si i-layer deposition
conditions.

  ---------------------------------- ----------------- ------------------
  Parameters                         HWCVD             PECVD

  Substrate temperature (°C)         500               175

  Pressure (mTorr)                   100               700

  SiH~4~ flow (sccm)                 9.3               2

  H~2~ flow (sccm)                   25                180

  Wire temperature (°C)              1850              \-\-\--

  RF power (W)                       \-\-\--           50

  Deposition time (min)              16                270

  Layer thickness (μm)               1                 1
  ---------------------------------- ----------------- ------------------

N/M=not measured

[]{#__RefHeading___Toc469989121 .anchor}Table 27. Undoped a-Si buffer
and Si p-layer deposition conditions.

  --------------------------- ------------------------ -------------------
  Parameters                  PECVD a-Si buffer layer  PECVD p-Si layer

  Substrate temperature (°C)  150                      150

  Pressure (mTorr)            200                      1000

  SiH~4~ flow (sccm)          20                       2

  H~2~ flow (sccm)            0                        190

  B~2~H~6~ flow (sccm)        0                        10

  RF power (W)                7                        50

  Deposition time (sec)       108                      240

  Layer thickness (nm)        30                       15
  --------------------------- ------------------------ -------------------

[]{#__RefHeading___Toc469989122 .anchor}Table 28. Device structures and
JV data.

+-----------+--------------------+-------+--------+--------+---------+
| Piece     | Structure          | W     | N      | Diode  | J~o~    |
|           |                    | (µm)  | (c     |        | (mA     |
|           |                    |       | m-^3^) | Factor | /cm^2^) |
|           |                    |       |        | A      |         |
+-----------+--------------------+-------+--------+--------+---------+
| G         | n+ Si/p+ µc-Si     | 0.03  | 2E18   | 2.1    | 2E-3    |
| D41015-21 |                    |       |        |        |         |
+-----------+--------------------+-------+--------+--------+---------+
| G         | n+ Si/a-Si/p+      | 0.04  | N/M    | N/M    | N/M     |
| D41019-21 | µc-Si              |       |        |        |         |
+-----------+--------------------+-------+--------+--------+---------+
| HW203-21  | n+ Si/HW i/p+      | 0.2   | N/M    | N/M    | N/M     |
|           | µc-Si              |       |        |        |         |
+-----------+--------------------+-------+--------+--------+---------+
| HW205-22  | n+ Si/HW i/a-Si/p+ | 0.2   | 2E16   | N/M    | N/M     |
|           | µc-Si              |       |        |        |         |
+-----------+--------------------+-------+--------+--------+---------+
| G         | n+ Si/GD i/p+      | 0.5   | 4E16   | 2.9    | 2E-3    |
| D41014-11 | µc-Si              |       |        |        |         |
+-----------+--------------------+-------+--------+--------+---------+
| G         | n+ Si/GD i/a-Si/p+ | 0.8   | N/M    | 4.6    | 5E-2    |
| D41018-21 | µc-Si              |       |        |        |         |
+-----------+--------------------+-------+--------+--------+---------+

The best devices were on pieces that consisted of only n+/p+ with no
µc-Si HW or GD absorbers. The JV results for these two devices are in
Table 29 and the JV curves are shown in Figure 71. The a-Si buffer
improves the V~OC~ but reduces the J~SC~ and FF. The increased current
in reverse bias in light and dark for the device without a buffer,
GD41014-21, is expected. It is due to a leaky n+/p+ tunnel junction. The
thin a-Si buffer layer eliminates this high tunneling current. The
highly doped n+ c-Si substrate (0.01 Ω-cm) is not expected to produce a
high J~SC~ or V~OC~ due to the short diffusion length.

[]{#__RefHeading___Toc469989123 .anchor}Table 29. JV results from n+-p+
cells having no GD or HW absorber.

  -------------- -------------------- ---------- --------------- ---------
  Piece          Structure            V~oc~ (V)  J~sc~           FF(%)
                                                 (mA/cm^2^)      

  GD41015-21     n+ Si/p+ µc-Si       0.39       11              59

  GD41019-21     n+ Si/a-Si/p+ µc-Si  0.62       2               37
  -------------- -------------------- ---------- --------------- ---------

![](media/image84.wmf){width="3.625in" height="3.7222222222222223in"}

[]{#__RefHeading___Toc469989202 .anchor}Figure 71. Light and dark JV
curves for n+-p+ cells GD41015-21

and GD41019-21 showing effect of a-Si buffer.

The QE measured in the dark at 0 V is shown in Figure 72. The peaks are
due to interference fringing. The n+-p+ cell (GD41015.21) response
extends to 1000 nm, while the others only extend to 700-800 nm,
indicating they have no photocurrent contribution from the n+ c-Si
substrate. The two devices with HW or GD i-layers with buffers had
identical QE curves. The GD device without the buffer had much higher QE
than the one with a buffer. The capacitance was measured at 100 kHz in
the dark. Results are given in Table 28. The n+-p+ cells have very small
depletion widths as expected since they are highly doped on both sides.
The space charge density for GD41015-21 of 2E18 cm-3 is as expected for
the doping of a n-type Si wafer with a resistivity \~1E-2 Ω-cm. The HW
and GD devices had comparable values of 2-4E16 cm-3.

In summary, the preliminary fabrication of n-i-p and n+-p+ cells using
GD and HW µc-Si absorber i-layers has produced cells with poor
photo-carrier collection and high recombination. The a-Si buffer layer
increases V~OC~ but reduces J~SC~ and FF in both n-i-p and n+-p+ cells.
No first-order differences were seen between devices the HW or GD µc-Si
layers. The highly doped µc-Si substrate does not appear to contribute
to the photo-carrier collection in the n-i-p cells. n+-p+ cells with the
a-Si buffer achieved V~OC~=0.62 V but with very poor J~SC~ and FF. The
following section discusses impurity characterization of these devices
in an attempt to explain their poor performance.

![](media/image85.wmf){width="3.5972222222222223in"
height="3.486111111111111in"}

[]{#__RefHeading___Toc469989203 .anchor}Figure 72. QE at 0V, dark for 4
devices.

### 4.2.3 Impurities in PECVD and HW films and devices

In the past two years, a decrease in FF has been measured from p-i-n
devices deposited by PECVD. Previously, devices routinely achieved
FF=70% while recently FF has been limited to 50-60%. Other device
parameters were comparable (V~OC~=0.86 V, J~sc~=16 mA/cm^2^). J-V
analysis indicated the low FF was not due to series resistance. The low
FF was not influenced by the back contact and brands of SnO~2~. QE
measurements were made as a function of bias voltage to help identify
the source of the low FF. Figure 73 shows the ratio of QE at 0.5V to 0V
for 2 devices, one with normal FF and a recent one with low FF. This
shape for the QE ratio of the low FF device is a well-known indication
of poor hole collection from the i-layer; i.e. greater bias dependence
at long wavelengths. The most likely cause is donor-like impurities such
as P or O, which reduce the hole lifetime. The sub-bandgap QE was
measured to characterize deep defects. Figure 74 shows log QE vs energy
for the same two devices. The device with low FF had higher absorption
at 1-1.2 eV confirming higher midgap defects.

![](media/image86.jpeg){width="3.7506944444444446in"
height="3.7506944444444446in"}

[]{#__RefHeading___Toc469989204 .anchor}Figure 73. QE ratios comparing
forward bias to 0V on two devices

showing greater bias dependence at long wavelength for device with low
FF.

![](media/image87.jpeg){width="3.7506944444444446in"
height="3.7506944444444446in"}

[]{#__RefHeading___Toc469989205 .anchor}Figure 74. Log QE plots showing
increased midgap

defects in device with lower FF.

The air leak-up rate into the PECVD reactor chamber was typically as low
as in previous years, indicating that an external leak was not causing
the problem. Considerable effort was made in cleaning the PECVD
deposition chamber to eliminate deposits, which could be a source of
residual O or P. Several other sources of impurities have been replaced.
The silane cylinder was replaced with a new cylinder, which was assayed
for purity. The internal components were replaced, made of Macor, which
insulate the electrodes and substrate holder. The IR heater lamps were
replaced. None of these changes resulted in a higher FF.

Select samples have been sent to NREL for SIMS over past 3 years.
Results are shown in Table 30 for both a-Si and µc-Si films and devices.
The first group are a-Si by PECVD, second is µc-Si p-i-n by PECVD, third
group are HWCVD i-layer and p-i-n cells, and fourth group is c-Si wafer
as control. Piece 4767 was a-Si p-i-n device with FF=70%. Piece 4997 was
a-Si p-i-n device with FF=55%. Silane cylinder and Macor internals were
replaced between 4976 and 4986. Not all elements were analyzed in every
sample. Values are for the bulk i-layer distribution only, except for
HW203 and HW205 where range is shown, due to significant tailing for C,
O, N from the free surface. PECVD a-Si films and devices, and µc-Si film
and devices have \~0.1 µm tail while HW µc-Si films have \>1 µm tail,
suggesting absorption of atmosphere in HW films as reported by several
other groups. Most HW films sent for SIMS in 1998 and 2001 were 5-10 µm
thick and also had \~1 µm tail for C, O, N but not other impurities.
Since recent devices were only 1 µm as discussed in the previous
section, the tail dominates the profile.

Since our PECVD reactor is a single chamber system, a major concern in
both a-Si and µc-Si devices is the P and B cross-contamination. Our
standard practice of depositing a contaminant burying SiC layer after
all device runs containing an n-layer was developed years ago in
response to serious P contamination. Table 30 shows that P and B are
quite low in PECVD a-Si films and devices. P was of critical importance
since Figure 73 showed that the low FF was due to a donor-like impurity.
The P level in the i-layers are near those for undoped c-Si wafers. They
are shown as "\<10^17^ cm^-3^" since there was considerable noise in the
SIMS profiles at these low levels. Note that several recent devices have
had P concentrations higher than 10^19^ cm^-3^ in the p-layer. 4989 and
4997 had B concentrations of 10^18^ cm^-3^ in the p-layer while P was 10
times higher. This suggests our burying layer is not effective anymore
in preventing cross-contamination.

A possible explanation is as follows. It has been shown by several
groups including ours that H-diluted i-layer depositions are much more
efficient at scavenging P compared to undiluted SiH~4~ plasma. New
p-layer conditions have required increasing the B~2~H~6~/H~2~ flow by
factor of 10 since the B~2~H~6~ source is diluted 0.2% instead of 2% as
used in the past. Thus, the H~2~/SiH~4~ ratio in the p-layer has
increased from 1.5/30 to 15/30, a tenfold increase in H dilution. This
may be responsible for increased P counter-doping in the p-layer.
Interestingly, the P cross-contamination does not seem to reduce the
V~oc,~ as might be expected nor does P appear to contaminate the
i-layer. Consequently, a new burying procedure has been implemented
using an initial H plasma step to scavenge P, followed by the standard
SiC step. Diagnostic device samples have been deposited following an
n-layer deposition and the new burying procedure. Device and SIMS
results will be reported.

[]{#__RefHeading___Toc469989124 .anchor}Table 30. Summary of SIMS for
IEC Si layers and devices.

+----------+--------+--------+-------+-------+-------+-------+--------+
| Sample   | Depo.  | SIMS   | O     | C     | N     | B     | P      |
|          |        |        |       |       |       |       |        |
|          | date   | Date   | (at/c | (at/c | (at/c | (at/c | (at/   |
|          |        |        | m^3^) | m^3^) | m^3^) | m^3^) | cm^3^) |
+----------+--------+--------+-------+-------+-------+-------+--------+
| 4764-11  | 02     | 05     | 6x1   | 2x1   | 3x1   | N/A   | \<1x   |
|          | /06/98 | /23/01 | 0^18^ | 0^18^ | 0^17^ |       | 10^17^ |
| a-Si     |        |        |       |       |       |       |        |
| layer    |        |        |       |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| 4767-12  | 02     | 05     | 6x1   | 3x1   | 3x1   | N/A   | \<1x   |
|          | /10/98 | /23/01 | 0^18^ | 0^18^ | 0^17^ |       | 10^17^ |
| a-Si pin |        |        |       |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| 4976-22  | 05     | 05     | 8x10  | 2x1   | 2x1   | N/A   | N/A    |
|          | /01/01 | /23/01 | ^18^- | 0^18^ | 0^17^ |       |        |
| a-Si     |        |        |       |       |       |       |        |
| layer    |        |        | 3x1   |       |       |       |        |
|          |        |        | 0^19^ |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| 4986-22  | 07     | 08     | 8x10  | N/A   | 1x1   | N/A   | N/A    |
|          | /31/01 | /10/01 | ^18^- |       | 0^17^ |       |        |
| a-Si     |        |        |       |       |       |       |        |
| layer    |        |        | 2x1   |       |       |       |        |
|          |        |        | 0^19^ |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| 4989-11  | 09     | 09     | 1x10  | 1x1   | 1x1   | \<1x1 | \<1x10 |
|          | /05/01 | /28/01 | ^19^- | 0^18^ | 0^17^ | 0^17^ | ^17^\* |
| a-Si pin |        |        |       |       |       |       |        |
|          |        |        | 2x1   |       |       |       |        |
|          |        |        | 0^19^ |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| 4997-21  | 04     | 01     | 4x1   | 7x1   | \<1x1 | \<1x1 | \<1x10 |
|          | /25/02 | /14/03 | 0^18^ | 0^17^ | 0^17^ | 0^17^ | ^17^\* |
| a-Si pin |        |        |       |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| 41018-21 | 12/02  | 01     | 5x1   | 3x1   | 8x1   | \<1x1 | \<1x   |
|          |        | /14/03 | 0^19^ | 0^19^ | 0^17^ | 0^17^ | 10^17^ |
| µc-Si    |        |        |       |       |       |       |        |
| pin      |        |        |       |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| HW181-11 | 0      | 04     | 2x1   | 7x1   | 4x1   | \<1x1 | N/A    |
|          | 3/9/01 | /19/01 | 0^20^ | 0^18^ | 0^17^ | 0^17^ |        |
| µc-Si    |        |        |       |       |       |       |        |
| layer    |        |        |       |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| HW203-21 | 12/02  | 01     | 4x10  | 4x10  | 8x10  | \<1x1 | \<1x   |
|          |        | /14/03 | ^20^- | ^20^- | ^18^- | 0^17^ | 10^17^ |
| µc-Si    |        |        |       |       |       |       |        |
| pin      |        |        | 2x1   | 1x1   | 2x1   |       |        |
|          |        |        | 0^19^ | 0^19^ | 0^18^ |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| HW205-22 | 12/02  | 01     | 4x10  | 3x10  | 1x10  | \<1x1 | \<1x   |
|          |        | /04/03 | ^20^- | ^20^- | ^19^- | 0^17^ | 10^17^ |
| µc-Si    |        |        |       |       |       |       |        |
| pin      |        |        | 3x1   | 1x1   | 2x1   |       |        |
|          |        |        | 0^19^ | 0^19^ | 0^18^ |       |        |
| (a-Si    |        |        |       |       |       |       |        |
| buffer)  |        |        |       |       |       |       |        |
+----------+--------+--------+-------+-------+-------+-------+--------+
| c-Si     | N/A    | 4      | 1x1   | 5x1   | 4x1   | 2x1   | 2x     |
| wafer    |        | /19/01 | 0^18^ | 0^17^ | 0^16^ | 0^14^ | 10^16^ |
+----------+--------+--------+-------+-------+-------+-------+--------+

> **but the P concentration was \~1x10^19^ in the p-layer. N/A = not
> available**

For µc-Si device samples 41018-21, HW203-21, and HW205-22, the
concentration levels of atmospheric impurities (O, N and C) are several
orders of magnitude higher than those in the a-Si cell and the c-Si
wafer. These impurities in the HW samples have a concentration gradient,
which spans the thickness of the i-layer and decreases from the top
surface. Prior SIMS analysis of HWCVD films with 5µm thickness
(HW181-11) has shown a similar gradient of these impurities, which level
off to a constant bulk concentration at approximately one micron into
the film. Similar gradients have also been reported in the literature
for HWCVD µc-Si films. In contrast, bulk concentration levels of the
same impurities in PECVD a-Si and µc-Si films are reached after 0.1 µm
from the surface. Impurities in the PECVD µc-Si seem to be uniform
suggesting they occur during growth. It is not clear at this point where
these impurities originate. However, it is important to remember that
both PECVD and HWCVD i-layers are exposed to the atmosphere prior to the
emitter layer deposition. In addition, it is speculated that the lower H
content in HWCVD films, combined with the excessive number of grain
boundaries and film porosity may facilitate contaminant penetration into
the films. To test this hypothesis, µc-Si i-layers will be capped with
a-Si:H and submitted for SIMS analysis. This will provide a baseline in
contaminant levels to be expected from both the PECVD and HWCVD
reactors. The control of oxygen incorporation into HWCVD films is a
critical issue as reports in the literature indicate several orders of
magnitude increase in the film conductivities as the oxygen content in
the film increases [^74] [^75] [^76].

### 4.2.4 Boron doped HWCVD Si films

In order to develop complete HWCVD Si devices, work was initiated on the
fabrication of doped Si layers. Boron doped HWCVD films were deposited
and characterized. The deposition conditions used are summarized in
Table 31.

[]{#__RefHeading___Toc469989125 .anchor}Table 31. B-doped Si film
deposition conditions.

  ----------------------------------------------- -----------------------
  Substrate temperature (°C)                      450

  Wire temperature (°C)                           1850

  Pressure (mTorr)                                200

  SiH~4~ flow (sccm)                              4.0

  B~2~H~6~/SiH~4~                                 10^-5^ -- 10^-3^

  Deposition time (min)                           30
  ----------------------------------------------- -----------------------

In addition, a set of undoped films was also deposited for comparison.
For these films, the deposition conditions were similar as in Table 31,
except that an additional 2 sccm of H~2~ was added to the gas mixture to
make up for the H~2~ diluent in the B~2~H~6~ dopant mixture. The film
thickness for all films varied from 0.56 to 0.73 m. XRD and Raman
spectroscopy analysis, as seen in Figures 75 and 76 for a representative
sample, shows that the films are highly crystalline and exhibit typical
(220) preferential orientation. Calculation of the crystalline fraction
from the Raman profiles results in values from 75 to 90%.

![](media/image88.png){width="3.779861111111111in"
height="3.6534722222222222in"}

[]{#__RefHeading___Toc469989206 .anchor}Figure 75. XRD pattern of
B-doped HWCVD μc-Si film.

![](media/image89.png){width="4.267361111111111in"
height="3.6430555555555557in"}

[]{#__RefHeading___Toc469989207 .anchor}Figure 76. Raman spectra of
B-doped HWCVD μc-Si film.

Values of the room temperature light and dark conductivity measured by
the 4-point probe method as a function of the B~2~H~6~/SiH~4~ in the gas
phase are summarized in Table 32. The invariance of conductivity with
dopant gas concentration indicates poor incorporation of B from the gas
phase at the conditions used. Previous experiments carried out at P =
500 mTorr and B~2~H~6~/SiH~4~ = 1x10^-3^ with other parameters similar
to those in Table 24 yielded films with conductivities of \~ 10^-2^
S/cm. However, the films had a significant thickness gradient along the
direction of flow due to the high pressure. A pressure of 200 mTorr was
used for this study to achieve greater uniformity. Although an extensive
study of B-doping by B~2~H~6~ has not been carried out, it is difficult
to rationalize how changes in pressure would affect the incorporation of
B into the films. Alternatively, it is possible that a change in
microstructure across the thickness of the films leads to differences in
the overall conductivity of the films. The samples deposited at P = 500
mTorr had a thickness of 4 m, a factor of 8 greater in thickness
compared to the films in this study. To determine the factors affecting
conductivity, a second set of B-doped films was deposited and their
properties compared to those of the samples in Table 32. Deposition
conditions and film properties for this sample set are given in Table
33. Comparison of sample HW223-11 in Table 33 and the samples in Table
32 clearly indicates that the deposition pressure does not play a role
in affecting the film conductivity. For submicron films, the
conductivity changes 3 orders of magnitude by changing the
B~2~H~6~/SiH~4~ from 10^-3^ to 10^-2^ as seen between sample HW226-11,
and the samples in Table 32. Furthermore, if the film thickness is
increased by a factor of 3 while keeping other conditions constant (see
samples HW226-11 and HW227-11), the conductivity increases another 2
orders of magnitude. The fact that the activation energy of both of
these samples is the same, indicates an increase in mobility with
thickness. The reasons for this dependence are not clear, however, a
possible explanation is the existence of a highly defective or amorphous
incubation layer in the initial stages of growth. Separating the effects
of dopant gas concentration and film microstructure on the film
conductivity will be an area of focus in future work.

[]{#__RefHeading___Toc469989126 .anchor}Table 32. Conductivity of
B-doped Si films deposited at 200 mT.

  ---------------- ----------------- -------------------- ---------------------
  Sample \#        B~2~H~6~/SiH~4~   Dark Cond. (S/cm)    Photo Cond. (S/cm)

  HW212-21         0                 1.1x10^-6^           1.6x10^-6^

  HW215-21         1x10^-5^          3.5x10^-6^           2.0x10^-6^

  HW216-21         1x10^-4^          2.4x10^-6^           2.3x10^-6^

  HW217-21         1x10^-3^          1.7x10^-6^           1.8x10^-6^
  ---------------- ----------------- -------------------- ---------------------

[]{#__RefHeading___Toc469989127 .anchor}Table 33. Deposition conditions
and film properties of B-doped Si films.

+----------+----------+---------------------+-------------+-----------+
| Sample   | T        | Deposition          | C           | A         |
| \#       | hickness | Conditions          | onductivity | ctivation |
|          |          |                     |             |           |
|          | (m)     |                     | (S/cm)      | Energy    |
|          |          |                     |             |           |
|          |          |                     |             | (eV)\*    |
+----------+----------+---------------------+-------------+-----------+
| HW120-22 | 4.0      | P =500 mTorr        | Dark=2.7E-2 | 0.072     |
|          |          |                     |             |           |
|          |          | B~2~H~6~/SiH~4~ =   | Photo=same  |           |
|          |          | 1.7x10^-3^          |             |           |
|          |          |                     |             |           |
|          |          | F~SiH4~ = 12 sccm   |             |           |
|          |          |                     |             |           |
|          |          | T~f~ = 1815 ^o^C    |             |           |
|          |          |                     |             |           |
|          |          | T~s~ = 450 ^o^C     |             |           |
|          |          |                     |             |           |
|          |          | Gr = 3.3 m/hr      |             |           |
+----------+----------+---------------------+-------------+-----------+
| HW223-11 | 1.5      | P =500 mTorr        | Dark=1.E-6  | 0.088     |
|          |          |                     |             |           |
|          |          | B~2~H~6~/SiH~4~ =   | Photo=same  |           |
|          |          | 1.0x10^-3^          |             |           |
|          |          |                     |             |           |
|          |          | F~SiH4~ = 4 sccm    |             |           |
|          |          |                     |             |           |
|          |          | T~f~ = 1850 ^o^C    |             |           |
|          |          |                     |             |           |
|          |          | T~s~ = 450 ^o^C     |             |           |
|          |          |                     |             |           |
|          |          | Gr = 3.0 m/hr      |             |           |
+----------+----------+---------------------+-------------+-----------+
| HW226-11 | 0.8      | P =200 mTorr        | Dark=1.5E-3 | 0.041     |
|          |          |                     |             |           |
|          |          | B~2~H~6~/SiH~4~ =   | Photo=same  |           |
|          |          | 1.0x10^-2^          |             |           |
|          |          |                     |             |           |
|          |          | F~SiH4~ = 3 sccm    |             |           |
|          |          |                     |             |           |
|          |          | T~f~ = 1850 ^o^C    |             |           |
|          |          |                     |             |           |
|          |          | T~s~ = 450 ^o^C     |             |           |
|          |          |                     |             |           |
|          |          | Gr = 1.2 m/hr      |             |           |
+----------+----------+---------------------+-------------+-----------+
| HW227-11 | 2.5      | P =200 mTorr        | Dark=1.1E-1 | 0.040     |
|          |          |                     |             |           |
|          |          | B~2~H~6~/SiH~4~ =   | Photo=same  |           |
|          |          | 1.0x10^-2^          |             |           |
|          |          |                     |             |           |
|          |          | F~SiH4~ = 3 sccm    |             |           |
|          |          |                     |             |           |
|          |          | T~f~ = 1850 ^o^C    |             |           |
|          |          |                     |             |           |
|          |          | T~s~ = 450 ^o^C     |             |           |
|          |          |                     |             |           |
|          |          | Gr = 1.2 m/hr      |             |           |
+----------+----------+---------------------+-------------+-----------+

**\*all films had curved Arrhenius plots over 213 to 373K, Ea from fit
only 213-293 K**

## 4.2.5 National Thin Si Team Activities

Robert Birkmire attended the National Thin Film Si Team Meeting held on
January 10. 2003. Bob gave a presentation on the growth and
characterization of HWCVD Si films on Al coated glass. Rogert Aparicio
attended the National Thin Film Si Team Meeting on August 8, 2003. Roger
gave a presentation on the grain enhancement of Si films by aluminum
induced crystallization approaches.

# **DIAGNOSTICS**

## 5.1. Contact Wetting Angle

The use of contact wetting angle between a film surface and a liquid is
being investigated as a diagnostic tool for thin-film photovoltaic
fabrication. The terminating surface of a solid contains excess free
energy compared to the bulk, arising in part from greater molecular
disorder where the bulk lattice is terminated. This excess free energy
is not easily measured by direct means but is manifest in the contact
angle of liquids in contact with the surface. The free energy of the
solid in air (SV) is related to the contact angle () and the free
energies of the solid to liquid (SL) and the liquid to air (LV) through
Young's equation [^77]:

$E_{\text{SV}} = E_{\text{SL}} + E_{\text{LV}}\text{cos}(q)$ (13)

By employing fluids with different viscosity and surface tension, the
equation can be solved simultaneously for the free energy of the solid
to air. For specular materials, performing these measurements on a
tilted plane reveals contact angle hysteresis between the leading and
trailing edges of the droplet, which provides quantitative information
relating to surface chemistry [^78]. Application of a voltage across the
droplet and the solid allows the contact angle to be modified; such
electro-wetting manipulation of the liquid allows determination of
interfacial free energy and hence the solid to air free energy [^79].
These techniques are most valuable when used to measure relative
differences in similar specimens.

Contact wetting angle measurements of water on CdTe films at different
stages of processing are correlated with surface phase composition.
Polar and dispersive components of the surface energy on Cu(InGa)Se~2~
films with different composition are determined. Reduction techniques
were implemented to give surface adhesive energy of the liquid to the
film and surface energy of the film.

For measurement of the wettability of liquid on a solid surface, a
useful comparative quantity is the adhesive energy, *E~ad~*, expressed
in terms of the liquids surface tension and the contact angle:

$E_{\text{ad}} = E_{\text{LV}}(1 + \text{cos}\theta)$ (14)

As the wettability of a liquid on a solid surface *decreases*, the
contact angle increases, ** → 90 *cos()*→ 0, and *E~ad~* → *E~LV~*.
In this case, the adhesive energy approaches the surface tension at the
liquid-vapor interface. As the wettability further decreases, **  90
and *cos()* \< 0. In this case the adhesive energy is less than the
liquid-vapor energy and the droplet will be poorly adhered to the
surface. As the wettability *increases*, the contact angle decreases,
** → 0 *cos()*→ 1, and *E~ad~*  *E~LV~*.

Since the solid-vapor energy is a fixed property, decreasing contact
angle implies a decrease in the solid-liquid interface energy, *E~SL~*→
0. The Zisman method of surface energy determination relies on
measurement of the contact angle using different fluids and plotting
*cos()* versus *E~LV~*. Extrapolation to *cos()* = 1 yields the
surface energy of the solid. Table 5.1 lists the properties of contact
liquids.

[]{#__RefHeading___Toc469989128 .anchor}Table 34. Properties of contact
liquids

+--------------------+-----------------+---------------+--------------+
| Name               | Formula         | T~b~          | *E~LV~*      |
|                    |                 |               |              |
|                    |                 | (°C)          | (dynes/cm)   |
+--------------------+-----------------+---------------+--------------+
| Acetophenone       | C~6~H~5~COCH~3~ | 203           | 39.8         |
+--------------------+-----------------+---------------+--------------+
| Ethylene Glycol    | HOCH~2~CH~2~OH  | 198           | 47.7         |
+--------------------+-----------------+---------------+--------------+
| Methylene Iodide   | CH~2~I~2~       | 182           | 50.8         |
+--------------------+-----------------+---------------+--------------+
| Formamide          | HCONH~2~        | 111           | 58.2         |
+--------------------+-----------------+---------------+--------------+
| Water              | H~2~O           | 100           | 73.0         |
+--------------------+-----------------+---------------+--------------+

The Zisman method was evaluated this year for Cu(InGa)Se~2~ and (CdZn)Te
alloys, with the result that the extrapolation method is too inaccurate
to achieve a reasonable degree of accuracy or repeatability. This is
because the method is best suited for liquids with low fraction of polar
energy. The total free surface energy consists of both polar and London
dispersive components:

$E = E^{p} + E^{d}$ (15)

The components are different for different fluids and can result in
non-linear behavior on a Zisman plot of *cos()* versus *E~LV~*. The
polar component, *E^d^*, arises from permanent dipoles, induced dipoles
and hydrogen bonding. The London dispersion force component of the
surface free energy, *E^p^*, arises from interactions between
fluctuating electronic dipoles with induced dipoles. A simple *rule of
thumb* is that a low energy liquid such as acetophenone will wet and
strongly adhere to a high energy solid such as stannic oxide. This rule
holds only if the intermolecular forces in both solid and liquid can
interact. The rule breaks down, however, when the the force interaction
is restricted, such as observed for water on oxide-free high energy
metals. In our studies, we intended to utilize water for routine
analysis. Thus, a more detailed reduction process was needed.

A more quantitative assessment can be obtained using two or more liquids
by the *harmonic mean* method in which the polar and dispersive
components of the solid, $E_{S}^{p}$ and $E_{S}^{d}$, are determined.
For liquid 1, we have $E_{1} = E_{1}^{d} + E_{1}^{p}$ and for liquid 2,
$E_{2} = E_{2}^{d} + E_{2}^{p}$:

$(1 + \text{cos}\theta_{1})E_{1} = 4\left( \frac{E_{1}^{d}E_{S}^{d}}{E_{1}^{d} + E_{S}^{d}} + \frac{E_{1}^{p}E_{S}^{p}}{E_{1}^{p} + E_{S}^{p}} \right)$
(16a)

$(1 + \text{cos}\theta_{2})E_{2} = 4\left( \frac{E_{2}^{d}E_{S}^{d}}{E_{2}^{d} + E_{S}^{d}} + \frac{E_{2}^{p}E_{S}^{p}}{E_{2}^{p} + E_{S}^{p}} \right)$
(16b)

Equations 16a and 16b can be solved simultaneously for the solid surface
energy components $E_{S}^{p}$ and $E_{S}^{d}$, given the contact angles
and corresponding polar and dispersive energies of the two fluids. The
polar fraction of the surface energy is represented as *x~p~*.

An apparatus for routine measurement of contact angle under standard
conditions was constructed and software for reducing wetting angle to
adhesive energy and solid surface energy was developed. Contact angle is
measured at room temperature (20-25°C) and atmospheric pressure using a
horizontal microscope at \~20X magnification, fitted with a protractor
reticle and *x-y-z* motion stage. A CCD camera is used to photograph the
drop at 20-100X as shown in Figure 77. Using the in-field protractor or
the on-screen protractor tool, contact angle can be determined with a
precision of ±1°. The fluids used for the present work were deionized
water and deionized formamide. For each sample and fluid, the contact
angle of the left and right side of 5 different drops were measured and
the average taken. Each drop was dispensed from a capillary pipette with
a volume of 0.05 cm^3^. Prior to measuring the contact angle, each drop
was allowed to reach equilibrium with the surface for \>10s. The samples
were rinsed in isopropanol (IPA) and dried under argon prior to contact
angle measurement.

![](media/image90.png){width="4.266666666666667in"
height="1.9319444444444445in"}

[]{#__RefHeading___Toc469989208 .anchor}Figure 77. Deionized water
droplet on the surface of VT CdTe.

Surface energy analysis has been carried out on Cu(InGa)Se~2~ films
deposited with similar Cu composition but different Ga content, all with
a final thermal step at 550°C (Table 5.2). The average contact angles of
water (*x~p~* = 0.32) and formamide (*x~p~* = 0.70) were determined for
the left and right sides of 3 drops. From the contact angle results, it
is apparent that water wets samples with Ga better than samples without,
while the wetting of formamide was unaffected by composition. Using the
Zisman analysis method, a surface energy of 55 dyn/cm was determined for
all samples. However, the results using the harmonic mean analysis
indicates a similar dispersive component for all samples but higher
polar component for samples with Ga. This preliminary result suggests
that the polar component of the surface energy in Cu(InGa)Se~2~ films
may be influenced by Ga concentration. A more detailed study of this
dependence is planned.

[]{#__RefHeading___Toc469989129 .anchor}Table 35. Contact wetting angle
and surface energy for Cu(InGa)Se~2~ films.

+---------+-------+-------+-------+-------+-------+-------+-------+
| Sample  | Cu/   | Ga/   | θ~w   | θ~    | E^p^  | E^d^  | E~SV~ |
|         | (III) | (III) | ater~ | forma |       |       |       |
|         |       |       |       | mide~ | (dy   | (dy   | (dy   |
|         |       |       |       |       | n/cm) | n/cm) | n/cm) |
+---------+-------+-------+-------+-------+-------+-------+-------+
| 3       | 0.82  | 0     | 43    | 14    | 29    | 31    | 60    |
| 3696.21 |       |       |       |       |       |       |       |
+---------+-------+-------+-------+-------+-------+-------+-------+
| 3       | 0.75  | 0.30  | 26    | 13    | 40    | 27    | 67    |
| 3690.21 |       |       |       |       |       |       |       |
+---------+-------+-------+-------+-------+-------+-------+-------+
| 3       | 0.77  | 0.31  | 18    | 14    | 45    | 25    | 70    |
| 3708.21 |       |       |       |       |       |       |       |
+---------+-------+-------+-------+-------+-------+-------+-------+

VT CdTe films supplied by First Solar as part of team activity were
analyzed by GIXRD and contact angle at different stages of
post-deposition processing. The GIXRD measurements were made on a Rigaku
D/Max 2500 instrument with CuKα x-rays at 40 kV and 50 mA, with incident
beam angle at 0.5°, sampling an average depth of \~60 nm. The contact
angle measurements of deionized water drops on the CdTe surface were
carried out at room temperature (20-25°C) and pressure using a
horizontal microscope at \~20X, fitted with a protractor reticle and
*x-y-z* motion stage. The contact angle was determined with a precision
of ±2°. For each sample, the contact angle of the left and right side of
6 different drops were measured and the average taken. Each drop has a
volume of \~0.05 cm^3^ and was dispensed from a capillary pipette
situated 2 cm above the level film specimen. Prior to measuring the
contact angle, each drop was allowed to reach equilibrium with the
surface for \>10 s. The films were rinsed in isopropanol (IPA) and dried
under argon prior to contact angle measurement.

The GIXRD and contact angle data are listed in Table 36. For as-received
VTCdTe films, only CdTe phase was found and the contact angle was
typically \>70°. Annealing the film in dry air at 600°C, with water
concentration less than 0.01 mg/L, caused insignificant change in
surface properties. Performing a CdCl~2~ vapor treatment at 430°C in dry
air produced CdO phase, detectable after 30 minutes of treatment, as has
been reported in previous studies [^80]. However, an increase in wetting
was detected after only 15 min. Similar adhesion was found for both 15
and 30 min treatments, suggesting, that the CdO phase is present in both
cases, but is thinner than the GIXRD detection limit for the 15 min
case. For CdCl~2~ treatment carried out in moist room air, with water
concentration \~0.01 g/L, the native oxides CdTeO~3~ and CdTe~2~O~5~
were found, and no CdO was detected. The proportion of CdTe~2~O~5~
increased with time, further reducing the adhesion. For the sample vapor
CdCl~2~ treated in dry air, the three-step BDH etch was used to produce
a Te layer, as shown by the GIXRD result. This significantly reduced the
wetting. After 30 min resting in room air, re-rinsed with IPA, and
analyzed, the wettting increased, suggesting the formation of Te oxides
too thin for GIXRD detection. This result is qualitatively similar to
that reported earlier for ellipsometric changes on Te surfaces produced
by NP etching [^81]. These relative measurements provide a sensitive
means to monitor changes in CdTe surface condition.

[]{#__RefHeading___Toc469989130 .anchor}Table 36. GIXRD and contact
angle results for CdTe surfaces at different stages of post-deposition
processing. Samples marked with (\*) are the same specimen.

+-------------+------------+----------------+-----------+-------------+
| Processing  | Conditions | Surface Phases | Contact   | Adhesive    |
|             |            |                | Angle     | Energy      |
| Stage       |            | from           |           |             |
|             |            |                | for H~2~O | for H~2~O   |
|             |            | GIXRD          |           |             |
|             |            |                | (±2       | (±2 dyn/cm) |
|             |            |                | degrees)  |             |
+-------------+------------+----------------+-----------+-------------+
| As          |            | CdTe           | 80        | 86          |
| -Received   |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
|             |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| Dry Air     | 600C, 20   | \+ CdTeO~3~    | 70        | 98          |
| Anneal      | min        |                |           |             |
+-------------+------------+----------------+-----------+-------------+
|             |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| CdCl~2~+Dry | 430C, 15   | CdTe           | 30        | 136         |
| Air         | min        |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| \*          | 430C, 30   | \+ CdO         | 25        | 140         |
| CdCl~2~+Dry | min        |                |           |             |
| Air         |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
|             |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| Cd          | 430C, 15   | \+ CdTeO~3~    | 35        | 132         |
| Cl~2~+Moist | min        |                |           |             |
| Air         |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| Cd          | 430C, 30   | \+ CdTeO~3~,   | 45        | 124         |
| Cl~2~+Moist | min        | CdTe~2~O~5~    |           |             |
| Air         |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| Cd          | 430C, 60   | \+ CdTeO~3~,   | 55        | 114         |
| Cl~2~+Moist | min        | CdTe~2~O~5~    |           |             |
| Air         |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
|             |            |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| \*BDH etch  | i          | Te             | 60        | 110         |
|             | mmediately |                |           |             |
+-------------+------------+----------------+-----------+-------------+
| \*BDH etch  | 0.5 hr     | Te             | 40        | 130         |
|             | after      |                |           |             |
+-------------+------------+----------------+-----------+-------------+

## 

## 5.2 Ellipsometry

Ellipsometry measures the complex p- and s- components of reflectance
from a specimen, in contrast to simply measuring the *intensity* ratio
of a reflected beam to the incident beam as in conventional
reflectometry. This is shown in the fundamental equation governing the
ellipsometry:

*tan Ψe^i∆^ =*$\frac{R^{p}}{R^{s}}$*,* (17)

where *R^p^* and *R^s^* are the total complex reflection coefficients
for the p- and s-waves, respectively. The parameter *Ψ* corresponds to
the magnitude ratio of the total reflection coefficient and *∆* is the
phase shift in p- and s- waves due to the reflection process:
Collectively they are known as the ellipsometry parameters. Ellipsometry
measures *Ψ* and *∆* and allows extraction of useful quantities such as
thickness and optical constants (n and k) in multi-layer structures
using detailed optical models. The knowledge of the polarization state
of the incident light as well as the reflected light are required for
the unique determination *Ψ* and *∆*. Thus an ellipsometer consists of a
light source, polarizer to fix the polarization of the incident light,
analyzer to measure the polarization state of the reflected light and a
detector. With these preliminary components, the ellipsometer can be
configured in a large number of modes; null, rotation analyzer, rotation
polarizer and polarization modulation.

Ellipsometry measurements were carried out on a J. A. Woollam VASE
system from 0.725 eV to 4.6 eV at intervals of 0.025 eV on single-phase
Cd~1-x~Zn~x~Te films deposited on CdS/ITO/glass substrate using
co-evaporation from CdTe and ZnTe sources. The VASE measurements were
carried out at room temperature under flowing Ar. A multilayer model was
developed to quantify the Cd~1-x~Zn~x~Te films in device structures, and
the optical constants were obtained from the best--fit. The accuracy of
the data and fit strategy was independently verified by comparing the
normal transmission data generated from the optical model with the
measured normal transmission. The optical data for Cd~1-x~Zn~x~Te films
is comparable to crystal data reported in the literature. The
compositional dependence of the various electronic transitions in the
Cd~1-x~Zn~x~Te films were fitted with second order polynomial equations,
and a semiconductor alloy model was developed using the optical
constants obtained from the Cd~1-x~Zn~x~Te films with different Zn alloy
composition. This alloy model has allowed the Zn alloy composition to be
accurately determined in unknown specimens in films prepared by IEC and
at the University of Toledo, deposited by PVD and rf sputtering.

## 

## 5.3 Glancing Incidence X-Ray Diffraction

X-ray diffraction (XRD) is a powerful yet non-destructive analysis tool
that provides quantitative information on the crystalline phases in thin
films and thin-film structures. With a suitable choice of x-ray source
and optical geometry, XRD analysis can be applied to a wide range of
polycrystalline and microcrystalline materials used for photovoltaics.
When used in conjunction with chemical analyses, a detailed description
of thin-film structures can be developed. For in-situ diagnostics, XRD
methods can provide a wealth of information on materials systems where
multiple phases are present or where the lattice parameter distribution
is coupled to deposition parameters.

The diffraction intensity pattern for monochromatic x-rays as a function
of the diffraction angle defined by 2*θ =* *θ~i~* + *θ~d~* contains two
primary pieces of information: 1) The overall pattern of diffracted peak
positions represents the instrument-transform of the family of lattice
plane spacing, which is relates to the specific crystalline phases
present; 2) Individual line profiles contain the instrument-transformed
distribution of d-spacings at that crystallographic \<*hkl*\> direction,
which in-turn is related to compositional or structural distortions of
the assemblage of atoms in that plane. Thus, XRD analysis consists of
decoupling instrumental effects (geometry and optical aberrations) from
specimen-related effects (phases, and properties thereof).

Two optical configurations that are routinely applied for thin film XRD
analysis derive from parafocusing cameras: *symmetric* Bragg-Brentano
geometry and *asymmetric* Seeman-Bohlin geometry. In either optical
configuration, the x-ray source intensity and angular dispersion are
fixed by the take-off angle from the x-ray source to the entrance slit.
The angular intensity distribution obtained on the detector arc
represents the convolution of the optical transformation function
(instrument function) with the specimen lattice distortion function. The
symmetric condition, *θ~i~* = *θ~d~*, constrains the measurement to
lattice planes parallel to the specimen plane, and the specimen is
equidistant from source and detector. The asymmetric condition, with
*θ~i~* ≠ *θ~d~* and using fixed incident beam angle, constrains the
measurement to populations of lattice planes that are tilted with
respect to the specimen plane. Asymmetric measurements made with low
incident beam angles are referred to as glancing incidence x-ray
diffraction (GIXRD). In either symmetric or asymmetric diffraction, the
minimum angle for diffraction from a specular specimen is governed by
the critical angle, *θ~c~*, for total external reflection, which for
CuKα radiation, is *θ~c~* \~ 0.2 for Si and *θ~c~* \~ 0.3 for CdTe and
CuInSe~2~. A more complete analysis of the evanescent and reflected
waves allows quantitative properties of multi-layer thin-film stacks to
be determined from glancing-incidence diffraction measurements.

The symmetric geometry is the most commonly encountered configuration in
commercially available XRD powder diffraction systems. The diffraction
depth depends on the mass absorption coefficient (μ/ρ), which in turn
depends on the x-ray energy and chemical composition of the phase,
causing low-index {hkl} planes (large d-spacing, low 2*θ*) to be
diffracted from near the powder surface, while high-index planes (low
d-spacing, high 2*θ*) are diffracted from deeper in the specimen.

Quantitative powder diffraction analysis relies on homogeneous
distribution and random orientation of the particulates, so that the
diffracted intensity distribution received at the detector is governed
only by the pure scattering properties of each {hkl} reflection of each
phase, and by the diffractometer instrument function. The detection
limit of a trace phase in a majority phase matrix depends on the
scattering properties of the phases; for similarly scattering phases,
the detection limit for symmetric XRD powder diffraction is \~0.5 wt%.
The measured volume fraction of surface layers can be enhanced by using
GIXRD measurements; utilizing the surface mass absorption coefficient to
constrain the sampling depth. Table 35 lists mass absorption
coefficients and sampling depths for different thin-film photovoltaic
absorber layers and Mo, Cu and Cr radiation for incident beam angle of 1
degree.

[]{#__RefHeading___Toc469989131 .anchor}Table 37. Mass absorption
coefficient (μ/ρ) and sampling depth (τ) for 1° incident beam angle and
different x-ray sources for thin-film photovoltaic absorber layer
materials.

  ------------------------------- ------------- ------------- -------------
  Mo Kα (0.017 MeV)               CdTe          CuInSe~2~     α-Si

  μ/ρ (cm^2^/g)                   35            69            8

  τ (μm) @ θ~i~ = 1°              0.85          0.44          11.5

  Cu Kα (0.008 MeV)               CdTe          CuInSe~2~     α-Si

  μ/ρ (cm^2^/g)                   260           130           50

  τ (μm) @ θ~i~ = 1°              0.12          0.23          1.8

  Cr Kα (0.0054 MeV)              CdTe          CuInSe~2~     α-Si

  μ/ρ (cm^2^/g)                   820           385           210

  τ (μm) @ θ~i~ = 1°              0.03          0.08          0.44
  ------------------------------- ------------- ------------- -------------

GIXRD analysis of crystalline surface phases in CdTe solar cells has
been correlated with device behavior for different post-deposition
processing conditions[^82]. Air heat treatment of CdTe films forms the
native oxide CdTeO~3~, while treatment in CdCl~2~/air vapor forms
CdTeO~3~ and CdO. The phases can exist as surface layers or, in some
instances, as penetrating phases. To first order, with the assumption
that the phases are randomly oriented, the presence of a surface phase,
A, on a substrate phase, B, can be determined by evaluation of the
intensity ratios for phases A and B at different incident beam angle.
The intensity ratios can be modeled using the phase scattering
properties and geometrical configuration. GIXRD measurements have
further shown that etching removes surface oxides and forms crystalline
Te. GIXRD is also being used to investigate changes in composition in
Cu(InGa)Se~2~ films, by compating the lattice parameter distributions
obtained at different incident beam angles, hence different sampling
depths.

In addition to phase information, GIXRD measurements are being routinely
used to determine through-film stress in as-deposited Mo films and in
CdTe films at different stages of processing. Analysis consists of
accurately determining the lattice parameter from GIXRD measurements
made at different incident beam angles spanning several (3-9) hkl
reflections. The shift in lattice parameters is transformed through the
strain coefficient tensor for the given material to give the residual
force in the plane of the film as a function of the depth. At present, a
series of sputtered Mo films deposited at fixed power but different
pressures have been fabricated and measured. Also, IEC PVD and VT
CdTe/CdS films and First Solar VT CdTe/CdS films have been measured
after different thermal processing steps and after surface etching. Full
analysis of the data is being carried out.

# **ABSTRACT**

This report describes results achieved during phase II of a three-phase
subcontract to develop and understand thin film solar cell technology
associated to CuInSe~2~ and related alloys, a-Si and its alloys and
CdTe. Modules based on all these thin films are promising candidates to
meet DOE long-range efficiency, reliability and manufacturing cost
goals. The critical issues being addressed under this program are
intended to provide the science and engineering basis for the
development of viable commercial processes and to improve module
performance. The generic research issues addressed are: 1) quantitative
analysis of processing steps to provide information for efficient
commercial scale equipment design and operation; 20 device
characterization relating the device performance to materials properties
and process conditions; 3) development of alloy materials with different
bandgaps to allow improved device structures for stability and
compatibility with module design; 4) development of improved
window/heterojunction layers and contacts to improve device performance
and reliability; and 5) evaluation of cell stability with respect to
illumination, temperature and ambient and with respect to device
structure and module encapsulation.

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