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

**Institute of Energy Conversion**

**United States Department of Energy**

**University Center of Excellence**

###### For Photovoltaic Research and Education 

Annual Report

to

National Renewable Energy Laboratory

under

Subcontract No. ADJ-1-30630-12

9/05/03-9/04/04

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-Based Solar Cells**

During this contract period substantial progress was made in
understanding and controlling the CdTe surface chemistry, in determining
the effect of humidity and pre-contact CdTe surface condition on
stability, in developing transparent ZnTe:Cu contacts, in establishing
reproducible CdTe deposition at high rate on a moving substrate in a
vapor transport reactor, in making devices with 12-13% efficiency on
commercial SnO~2~, and in developing new methods for device analysis.

Little is known regarding the effects atmospheric oxygen and humidity on
device performance and chemistry during standard stress conditions. We
have carried out a systematic investigation of the effects of stress in
the dark and light at 80°C at various humidities on IEC and FS CdTe
devices. The formation of TeO~2~ is observed at 0% humidity but no
TeO~2~ phases are observed at higher humidities. Cu~2-x~Te in the back
contact is observed at all humidities. The formation of a new phase,
CuTeO~3~, is observed in the pattern for the 0% humidity stressed
device. IEC cells exhibited forward-reverse retrace hysteresis prior to
stressing which became worse with stressing. The 0% humidity stressed
device showed the best stability despite the significant chemistry
occurring in back contact during stressing. No back contact changes were
detected by GIXRD for the FS stressed devices. All stressed FS devices
exhibit J-V rollover at forward bias and decreased V~oc~, increasing in
severity with higher stress humidity. The J-V data of the dark and
illuminated humid stressed devices shows similar trends between the two
stress conditions, though more rapid degradation is observed for the
devices stressed under illumination. We have attributed the improved
stability of the IEC devices to the thicker Te.

A common etchant used for CdTe device processing is the nitric
acid/phosphoric acid mixture (NP), which produces a 50-100 nm
crystalline Te layer on CdTe. Experiments were performed to study the
autocatalytic nature of NP etching. When using fresh NP solutions it was
observed that the etch rates of CdTe were considerably slower than
expected. Subsequent etches in the same solution showed faster etching
as evidenced by bubbling time, suggesting that a product of the reaction
is itself involved in the reaction. Bubbling times were found to
decrease with increasing NO~2~^-^, with almost immediate bubbling
observed with treatment of CdTe in NP containing 0.16M NaNO~2~. From
these results, the use of NO~2~^-^ based etchants for CdTe back
contacting appears promising and may benefit for CdTe device processing
by allowing shorter immersion times, more efficient and controllable
etching, and avoiding the use of concentrated HNO~3~.

The University of Toledo group has recently presented results regarding
a new chemical treatment for CdTe device processing, where samples are
illuminated while immersed in an aniline-based aqueous solution. IEC has
collaborated with the University of Toledo to assist in understanding of
the solution and surface chemistry that may be involved in the
aniline-based process. VASE and GIXRD both confirm the presence of
crystalline Te \~50-100 nm on the CdTe surface following successful
treatments under illumination. Results were very sensitive to the
illumination during the treatment. A successful treatment requires the
presence of aniline, Cl^-^ ions, O~2~, consistent illumination and
careful control of pH. Devices have been processed using aniline
treatments under illumination and completed with Cu-containing back
contacts and graphite paste. Initial devices show very promising
performance, with V~oc~ of ∼800 mV, J~sc~ ∼20 mA cm^-2^, and fill factor
∼60% or higher. We conclude that a successful aniline treatment of CdTe
results in the formation of a crystalline Te-rich surface layer on the
CdTe surface, which will be beneficial for back contact processing.

Development of ZnTe:Cu contacts has been motivated by the high p-type
conductivity of the ZnTe, needed for low contact resistance, and as a
source of Cu doping the CdTe, needed for high V~oc~. In addition, ZnTe
can serve as a semi-transparent back contact for bifacial JV and QE
measurements, to provide new insight on the device physics and back
junction properties, and as a transparent contact for the top cell of a
polycrystalline tandem. During the present contract period, IEC
investigated the galvanic solution growth of ZnTe. Tri-ethylamine (TEA)
was investigated as a complexing agent to control the chemical
reactivity of the Cu in the solution. Cells with optimum TEA had 10%
efficiency and negligible hysteresis. The light JV curve indicates a
weakly blocking contact when illuminated from front (through glass) but
when illuminated from back (through ZnTe) does not have a blocking
contact. This indicates that the back contact is photoconductive.
Backwall QE measurements through the transparent ZnTe:Cu/ITO shows the
response is largely limited to carriers generated very near the CdS
junction. There is no evidence of collection at a back junction, nor is
there evidence of much collection throughout the CdTe which implies a
low lifetime and low field in the bulk of the CdTe.

Increasing V~oc~ with the thin CdS needed for maintaining high Jsc
requires reducing the effect of pinholes or regions of non-uniform CdS.
The degree of spatial non-uniformity can be directly measured using a
modified scanning probe microscope (SPM). We have refined sample
preparation procedures to yield polished cross-sections of working cells
and have modified SPM for operation in different electrical
configurations. Electric field gradient (EFG) measurements have been
made with and without additional illumination and at different junction
bias. Measurements made at forward and reverse bias and at different
illumination conditions revealed marked changes in gradient and have yet
to be fully interpreted.

Several baseline VT devices were fabricated to compare with 60 nm thick
Ga~2~O~3~ and In~2~O~3~ high resistance transparent (HR) layers. CdS
films were deposited by chemical surface deposition to a thickness of
100 nm. Cells with a Ga~2~O~3~ HR layer had V~oc~ of \~0.79 V and
efficiency of \~12% while those with an In~2~O~3~ HR layer had V~oc~
\~0.60 V and efficiency of 8%. Therefore, Ga~2~O~3~ HR layers were used
in subsequent VT device studies.

Significant design modifications were made to the vapor transport source
to increase CdTe utilization and to improve repeatability and overall
robustness. The pre-deposition and post-deposition zone heaters were
replaced. The new heaters use Tantalum wire looped between a sheet of
Vycor glass and a grooved boron nitride plate. A new source was designed
and fabricated. The most significant change was the method of delivering
heat to the ampoule. Problems with the original design resulted in
contamination of the CdTe from the filament as well as significant CdTe
deposition onto the inside surfaces of the manifold. The CdTe
utilization at the substrate with this source had been 10-20%. The
contamination resulting from filament degradation, especially Si and Al,
caused shunting and poor efficiency and was verified by SIMS. The new
design utilizes a double-wall cylindrical boron nitride source heater to
heat the quartz ampoule containing the CdTe charge. Over the range of
experimental conditions, re-sublimation of the deposited film is not a
critical issue and the CdTe utilization is found to be 50% ± 5%. All
films exhibit faceted morphology, with means lateral grain size
proportional to film thickness. Grains with lateral dimension comparable
to film thickness, \~ 6 µm, are found at 550 and 570°C substrate
temperature, but the average grain size is higher and the distribution
of sizes is narrower at 570°C.

Alternative processes for treatment in Cl~2~ or HCl vapor have also been
examined. Two atmospheric pressure approaches for delivery of CdCl~2~
vapor species were considered: diffusion in parallel-plate configuration
and vapor transport via carrier gas in a packed bed. In both approaches,
the diffusion of CdCl~2~ through ambient gas is a critical aspect for
estimating concentration at the film surface. Although comparable
standard deviations were found for the wet and parallel plate vapor
treatments, the highest performance, with efficiency = 12%, V~oc~ \> 800
mV and FF \> 73%, was obtained with parallel-plate CdCl~2~ vapor
treatment. The cells made with packed bed CdCl~2~ treatment had lower
performance and uniformity.

The sensitivity of cell performance to VT deposition and post-deposition
processing, in particular the effect of CdTe growth rate and CdCl~2~
vapor treatment, was evaluated using CdTe from the newly modified VT
reactor. All samples were fabricated on TEC15 soda-lime glass/SnO~2~,
with 60 nm Ga~2~O~3~ HR layer and 80-90 nm CdS. The CdTe films were
deposited at 20 Torr using He carrier gas and an O~2~ base pressure of
\~ 0.02 Torr. There are four key conclusions. First, equivalent results
were obtained with and without a post-deposition anneal at T\>550°C in
argon, showing that the anneal step is unnecessary, unlike PVD and
electrodeposited cells, in which best results are obtained with an
annealing step prior to CdCl~2~ treatment. Second, cells with no CdCl~2~
treatment exhibit marked photo-response, indicating that the high
deposition temperature yields an active junction, again in contrast to
PVD or electrodeposition. Third, similar results were obtained over the
range of static equivalent growth rate, from 7 to 12 m/min and
deposition temperature, from 550ºC to 570ºC. Fourth, nominally similar
performance was obtained for vapor CdCl~2~ treatment from 405ºC to 425ºC
for 20-25 min but the highest V~oc~'s were obtained over a narrower
range, from 405ºC to 410ºC. Cells with no post-deposition Cl treatment
exhibited the lowest blue response (thickest CdS) and the lowest
absorber band edge (highest band gap), consistent with no CdS diffusion
into the CdTe. The other samples show a similar range of blue response
and similar long wavelength fall-off, consistent with CdS loss ≈ 10 nm
and formation of CdTe~1-x~S~x~ with x \< 0.05, respectively. Devices
with V~oc~\>0.80 V and FF\>68% were obtained on pieces from multiple
runs using Au/Cu evaporated contacts. The highest efficiency was on
piece VT118.4 with Voc=0.808V, Jsc=23.8 mA/cm2, FF=69.2%, and
efficiency=13.3%.

Voltage dependent photocurrent collection losses can be especially
significant in CdTe solar cells, primarily reducing the fill factor. We
analyzed the J~L~(V) effect in CdTe solar cells by measuring
current-voltage curves J(V) at different intensities and using the
difference to quantify J~L~(V) similar to previous work at IEC on CIGS.
We have analyzed over 10 CdTe devices from 4 different CdTe sources
having FF from 50 to 72 and thicknesses from 1.8 to 5 µm. All show
significant J~L~(V) losses affecting the maximum output power and FF.
Recalculating the FF without these losses results in a 5-15% point
increase. Voltage dependent collection in CdTe solar cells increases
after stressing. This effect is well described using the uniform field
model as was developed for *p-i-n* solar cells. Numerical simulations
with AMPS support the experimental results.

Measurements of V~oc~ (T) shows linear fit from 220 to 320K on most CdTe
devices over a range of intensity with an extrapolated intercept of
\~1.5 eV consistent with the CdTe bandgap. However, we have found that
at lower temperatures, the V~oc~ becomes independent of light intensity
and temperature, saturating at about V~oc~=1.0 V. We suggest that this
upper limit represents a fundamental parameter, the built-in voltage.
Saturation of V~oc~ indicates carrier freeze-out. When there are no
longer sufficient free carriers, the quasi-Fermi levels no longer
respond to changes in carrier generation or temperature. V~oc~ reaches
its maximum value determined by the absorber bandgap and the contact
potentials. The saturated value of V~oc~ decreases after stress from
\~1.0V to \~0.90-0.95 V.

***Team activities and collaboration***

The Institute has maintained a high level of collaboration with CdTe
Thin Film Partnership team members. Towards the goal of helping First
Solar to develop vapor CdCl~2~ treatments, IEC: constructed a new
CdCl~2~ source generator at IEC to carry out experiments compatible with
First Solar; carried out experiments to directly compare treatment in
CdCl~2~ vapor with treatment of CdCl~2~-coated samples; submitted
detailed reports on modeling CdCl~2~ vapor delivery; and investigated
source-poisoning issues associated with oxide formation in the source
generator. IEC performed GIXRD measurements on a large number of CdTe
films at different stages of processing and for different back contact
processing conditions. IEC characterized the structure and composition
of the back contact surface of First Solar production plates and of
newly developed contacts made by a third party company for First Solar.
IEC has also been assisting in the analysis of CdS film properties,
CdTe-CdS interaction after processing, and CdS/TCO quality assurance
issues for First Solar. IEC received a set of samples from First Solar
including baseline devices, as well as those having unusual J-V
performance. Complete analysis of J-V-T and QE(V) was performed and
reported at the team meeting along with results on similar samples by
CSU and USF.

Collaboration with Canrom has also been on-going, with respect to
post-deposition processing methods for Canrom CdTe/CdS and to the
Team-related issue of evaluation of humidity during cell stress. This
follow-up to the samples stressed in the dark at 85°C at different
humidities consist of repeating the stress conditions but in the light
with the calls at V~OC~. Results were presented at the team meeting and
in this report.

We also provided devices to CSM for cross-sectional luminescence
measurements.

Darshini Desai, Kevin Dobson, Brian McCandless and Steve Hegedus
attended the CdTe National Team Meeting in Toledo, February 2004. Much
of the work presented there is also contained in this report: the
chemistry of the aniline bath treatment of CdTe; the effects of humidity
stress of CdTe under illumination; two relatively new device
characterization techniques; and work related to the FS vapor-CdCl~2~
treatment.

## **CuInSe~2~-Based Solar Cells**

The process for the roll-to-roll deposition of Cu(InGa)Se~2~ on polymer
substrate has been improved during the past year by developing several
system improvements and reducing or eliminating the cracking of the Mo
back contact film in the Cu(InGa)Se~2~ reactor. The system was made more
robust and stable with a new evaporation source design and improved Se
delivery system. The new Cu source has a design, which avoids sharp
internal lines and corners that acted as stress concentrators. In
addition, the design includes conical nozzles that eliminated spitting
of Cu droplets and allows higher rate deposition. The Se delivery
manifold was simplified and reduced to a single 3/8\" stainless steel
tube with two effusion holes. A larger Se source was designed and
implemented which incorporates improved ability to measure and control
the melt temperature as well as greater capacity. Cracking of the Mo
film on polyimide substrates after Cu(InGa)Se~2~ deposition has been a
significant problem hampering progress with the roll-to-roll flexible
Cu(InGa)Se~2~. The extent of film cracking was reduced with lower excess
Se flux in the reactor implying that reaction of the Mo back-contact
with Se was responsible for crack formation by reducing the yield
strength of the Mo film. Since Mo oxides are more stable than the
selenides, Mo films were prepared containing around 8 at% oxygen by
sputtering in a mixed Ar/O~2~ atmosphere. It was thought that
introducing controlled oxygen content into the Mo film could reduce
reaction with Se and thus eliminate cracking. Based on this line of
logic. These films did not show any cracks, as evidenced by microscopic
characterization and by device performance, indicating that the
oxygenated Mo had reduced reaction with Se to prevent cracking. With
these changes to the system and substrate, improved deposition
reproducibility and compositional and device uniformity was
demonstrated.

Diffusion of sulfur into Cu(InGa)Se~2~ thin films could provide a means
to increase V~oc~ by widening the bandgap in the space charge region of
the absorber. A fundamental study of the mechanisms of S-diffusion was
completed using reaction of thin films and crystals in H~2~S. Two modes
of S incorporation in the CuInSe~2~ were distinguished. Rapid S
incorporation with a breakdown of the crystal lattice occurred in
materials containing excess Cu. In single-phase films or crystals, a
slow S incorporation was observed. Quantitative models of the
S-diffusion and diffusion coefficients were developed for these two
cases. Also, S was incorporated into evaporated CuInSe~2~ and
Cu(InGa)Se~2~ films by annealing them in diluted hydrogen sulfide
combined with oxygen and hydrogen selenide. After reaction in H~2~S or
H~2~S combined with O~2~, a fully converted sulfide layer was formed at
the surface of the absorber layer, which would create a barrier for
current collection in devices. A mixed hydride gas H~2~S + H~2~Se a
graded Cu(InGa)(SeS)~2~ layer on the top of the absorber layer was
formed and at the same time the Ga diffuses away from the surface. Thus,
the bandgap in the near surface region could not be controllably
increased with S diffusion so it does not appear to be a promising
avenue for increased V~oc~.

Research on CuInSe~2~-based devices with increased V~oc~ has focused on
determining the cause for poor adhesion which has limited the ability to
fabricate wide bandgap Cu(InAl)Se~2~ devices, and on S-diffusion in
Cu(InGa)Se~2~ for widening the bandgap in the near surface region. Poor
adhesion at the Mo/Cu(InAl)Se~2~ interface was identified as a critical
issue and was previously improved by incorporating a thin, \~ 5 nm, Ga
layer sputtered onto the Mo film prior to the Cu(InAl)Se~2~ deposition.
Still, devices with increasing bandgap or using different deposition
processes were severely hampered by poor Mo/Cu(InAl)Se~2~ adhesion. Poor
adhesion at the Mo interface can also be problematic for the fabrication
of Cu(InGa)Se~2~ cells on flexible polyimide substrates. Research is
focused on determining how Ga and/or Al at the Mo interface affects the
formation and orientation of a MoSe~2~ interlayer. The MoSe~2~ forms a
layered structure and orientation of these layers parallel to the
substrate can lead to poor adhesion. Mo, Mo/Cu and Mo/Ga layers were
reacted with Se and x-ray diffraction measurements showed that the
MoSe~2~ had a more preferred orientation with Ga than with Cu. Similar
experiments with Al could not be completed since it was limited by the
formation of Al~2~O~3~ or, with Mo/Cu/Al layers, formation of Cu-Al
intermetallic compounds. A second method was to characterize the MoSe~2~
after peeling off the CuInSe~2~-based films but the remaining layers
were found to be unstable.

Fundamental characterization of the optical properties of Cu(InGa)Se~2~
continued. Previously, a method to determine optical constants of
Cu(InGa)Se~2~ films using spectroscopic ellipsometry (SE) was developed
and variation with Ga/(Ga+In) ratios from 0 to 1 was determined. In this
phase, we have focused on the effect of Cu stoichiometry in Cu-poor
materials since high efficiency devices are made using single or
two-phase Cu-poor films. For both Cu-In-Se and Cu-In-Ga-Se alloy films,
an optical model was developed which allows the optical constants of
mixed phase materials to be described by a mixture of the single phases,
CuInSe~2~ and CuIn~3~Se~5~ for the ternary case. Complete determination
of the optical constants for films with Cu at. % from 11 to 25 % was
completed. This enables the composition of the mixed phase materials to
be determined using the SE measurements. For each alloy system, there is
an increase in the energy of the lowest order transitions, which
determine the bandgap. Also, there is a broadening of critical point
feature, which indicated decreasing crystalline quality.

The effect of the Cu(InGa)Se~2~/emitter layer interface and band
alignment on devices is being studied by changing the alloy compositions
of the Cu(InGa)Se~2~ layers and using CdS, (CdZn)S or ZnS emitter
layers. Also, chemical bath (CBD) and high utilization chemical surface
(CSD) depositions were compared. A process for CSD (CdZn)S was developed
and increased bandgap with Zn addition was confirmed by QE measurements.
In all cases, the (CdZn)S gave comparable performance to the CdS emitter
layer with both 1.2 and 1.4 -- 1.5 eV absorber layer bandgaps. The ZnS
emitter layers gave poor device efficiency. The CSD process resulted in
\~30 mV lower V~oc~ than the CBD process which was not improved with an
ammonia surface treatment prior to CdS deposition similar to the first
step in the CBD process, indicating the complexity of the interface
chemistry in controlling device behavior.

***Team activities and collaboration***

IEC activities in support of the CuInSe~2~ National Team and other
research groups working on CuInSe~2~ development have included the
following:

[Thin Film Partnership CuInSe~2~ National Team:]{.underline} IEC has
collaborated with several members of the Absorber Sub-Team to complete a
study comparing characterization of films and devices from 6 different
groups. In particular, IEC completed analysis of J-V and QE data on
devices it had previously fabricated using different Cu(InGa)(SeS)~2~
absorber layers.

[University of Illinois:]{.underline} Cu(InGa)Se~2~ samples were
provided to the U. of Illinois for TEM and other microstructural
analysis. Specific samples included films deposited at different
temperatures, comparing uniform flux, two-step, and three-step
processes, and deposited on Na-free glass for comparison to films grown
with Na diffusion from the soda lime glass substrate.

[University of Oregon:]{.underline} IEC continued its collaboration with
the University of Oregon in the analysis of Cu(InGa)Se~2~ based devices.
Specifically, IEC has provided various devices using different absorber
layers and substrates.

[Miasole:]{.underline} IEC fabricated solar cells and completed JV and
QE measurements to help Miasole characterize their Cu(InGa)Se~2~
deposition process.

[InterPhases Research:]{.underline} IEC completed cell fabrication and
provided J-V testing on a set of samples from InterPhases Research to
test their novel surface treatments. In addition, IEC has provided
glass/Mo/Cu(InGa)Se~2~ films for treatment

## **Si-Based Solar Cells**

Post-deposition annealing experiments of Al/Si bi-layers were carried
out to compare crystallization and layer inversion processes in standard
metal induced crystallization and in-situ metal induced crystallization.
Evaporated Si films allowed comparison between films deposited in an H
free environment and HWCVD films. First, 0.5 mm thick films of Al were
deposited onto 7059 glass substrates by electron beam evaporation. The
films were left in ambient air for a period of 24 hr to allow formation
of an oxide layer. Si films with 0.6 mm thickness were then deposited by
electron beam evaporation and HWCVD. XRD measurements showed no evidence
of crystallization of the Si films deposited by evaporation prior to
annealing while films deposited by HWCVD were completely crystallized.
Annealing was performed for 90 min at 500°C under 100 sccm of Ar. For Si
films deposited by evaporation, optical microscopy showed
polycrystalline Si nucleation extending to the glass interface with very
little lateral grain growth. XRD analysis showed only partial
crystallization of these films. For samples with Si films deposited by
HWCVD, no increase in grain coalescence or crystalline fraction was
observed.

Work continued on in-situ crystallization during HWCVD growth. Analysis
of HWCVD Si films deposited on Al coated glass substrates below the
Si-Al eutectic temperature (\~570°C) revealed a layered microstructure.
the lower layer near the glass interface is composed of micron size
islands of polycrystalline silicon. Atop this layer, a mixture of
nanocrystalline Si and Al remains. This mixture is also observed in the
spaces between the polycrystalline Si islands. The existence of this
layered structure is evidence of layer inversion and is consistent with
standard metal induced crystallization results obtained by other groups.
However, the discontinuity of the underlying polycrystalline Si
structure limits its usefulness for device fabrication. Raman
spectroscopy through the glass substrate confirms regions of
polycrystalline Si and regions of a mixture of nanocrystalline Si and
Al. After etching the overlaying Si-Al mixture, a similar pattern of
polycrystalline Si islands is observed from the front surface.
Post-deposition anneal at 430 °C under vacuum for 2 hours increased the
fractional coverage of c-Si from 60 to 90%. This indicates that the
excess Si does not crystallize as it accumulates on the Al layer surface
during deposition and is available for further lateral growth of the
polycrystalline Si layer.

The average grain size appears to decrease with increasing Si/Al ratio
and correlates well with the time the Al layers remained exposed to air
prior to Si film deposition. The role of aluminum oxides on the layer
inversion process and grain size of the polycrystalline Si layer is not
well understood. It is hypothesized that the oxide layer acts as a
nucleation barrier leading to a lower density of near neighbor nuclei.
However, it is expected that the oxide layer is self-limiting and would
reach equilibrium after a brief period of time of exposure to air. The
present results suggest that the oxide layer continues to change long
after deposition.

In order to consider the effect of growth rate and Si/Al ratio, Si films
with thickness of 2-3 μm was deposited by HWCVD at 5 μm/hr at 430°C on
Al layers 100, 250, 500 and 1000 nm thick. Optical micrographs taken
through the glass substrate of the Si films showed large areas of
polycrystalline Si grains. The central area of the film on 500 nm Al,
appears to have the densest coverage approaching 100%. The edge areas of
this film had a low fractional formation of Si grains. However, after
the 8 hr post deposition anneal at 500°C, the edge area had a similar
high fractional coverage as the middle. These continuous polycrystalline
layers appear as promising candidates for device fabrication.

It has been reported that the aluminum oxide layer is required for the
layer exchange process, although it is not required for the
crystallization process itself. It also acts as a diffusion barrier for
Si and Al atoms leading to a lower nucleation and thus large Si grains.
The growth of the oxide represents a potentially uncontrolled variable
in much of the MIC work. The oxide surface of Al films was monitored
over a period of 12 days using spectroscopic ellipsometry (SE)
measurements. Samples of e-beam deposited Al on glass were stored in
air. Although formation of the native oxide layer is thought to be a
self-limiting process, SE reveals changes in the optical properties of
the oxide layer over the 12 days indicating that its structure may be
changing long after the Al film deposition.

Previously, we have reported that such high temperature Si deposition on
Al was problematic since the Al films appear to become discontinuous.
Recently, it was found that Al films heated to 600°C in vacuum remained
continuous. Buffer layers between the glass and Al will be studied to
see if they can maintain wetting on the glass, hence, give continuous Al
films with Si film deposition at T\>600°C.

Regarding solar cell development, we have worked on solving dopant cross
contamination in our PECVD system and on developing B doped HWCVD layers
for use as an absorber for a thin film Si device. Recently, a-Si devices
fabricated in the single chamber PECVD reactor had lower fill factors
compared to devices fabricated several years ago. QE analysis indicates
that the lower FFs are caused by poor hole collection in the i-layer,
most likely arising from donor impurities such as P or O. We were
concerned about cross-contamination since the PECVD system will be
needed for deposition of thin a-Si n and i layers for HIT-type emitters
on the HWCVD p-type base layer. Several a-Si i-layers and devices were
sent to NREL for SIMS characterization. Concentrations of the
atmospheric contaminants N and C in the i-layer of all samples were low
and comparable in magnitude, and relatively uniform in distribution,
suggesting that air leaks are not responsible. For most samples, the P
concentrations were below the c-Si background level. But a recent
sample, having low FF, had double the O concentration compared to the
earlier ones. This sample was the only one to have measurable P
contamination.

Several PECVD device runs were made with varying burying layer, bake-out
procedures, and p-layer flows. Bias dependent QE measurements showed a
very clear dependence on the p-layer gas flow conditions. The data is
consistent with the longer p-layer contributing more B dopant
incorporation in the i-layer, which was not removed by our standard p-i
interface pump/purge cycle and a-SiC buffer layer, thus leading to poor
electron collection. Improved burying layers and reduced p-layer B flow
will be investigated to control and reduce cross-contamination.

The deposition of boron doped HWCVD Si films was investigated as part of
our effort to develop polycrystalline thin film Si devices. The films
were deposited at 500 m Torr and 450 °C with B~2~H~6~/SiH~4~ flow ratios
varying from 10^-5^ to 10^-3^. For all dopant gas concentrations, the
resulting dark conductivity of the films was 10^-6^ S/cm. In order to
understand whether the invariance in conductivity was the result of the
deposition conditions or film structure, doped films were deposited with
varying thickness. Increasing the thickness by a factor of 3 increases
the conductivity by 2 orders of magnitude but did not change the
activation energy, indicating an increase in mobility with thickness.

## **Diagnostics**

The in-line process diagnostic effort is intended to develop the tools
needed for process control and quantitative models to link sensor output
to process variables to material properties. Commercial-scale PV module
fabrication would benefit from methods to analyze thin-film surfaces at
various stages of processing. Contact wetting angle measurement is a
simple diagnostic tool that may be suitable in industrial applications
for monitoring changes in thin film solar cell processing conditions.

The terminating surface of a solid is known to contain excess free
energy when compared to the bulk. Surface energies (E^P^) consist of
polar energies (E^D^) resulting from permanent and induced dipole
moments as well as hydrogen bonding, and dispersive energies that arise
from vibrational interactions between atoms. The interaction of liquids
and solids of different surface energies are known to give rise to the
formation of characteristic angles known as wetting angles at a
liquid/solid interface.

Wetting angles were recorded for water droplets deposited onto CdTe
surfaces at varying times over the course of an hr immediately (\<1 min)
upon removal from the evacuated deposition chamber, during exposure to
air and after subsequent chemical etching for 5 sec. in a
bromine-methanol solution. A CdTe film tested upon removal from the
deposition chamber (as-dep) nearly completely wet the entire surface of
the sample (θ \~1^o^). Rapid increases in this value were observed over
the first hr. Bromine-methanol etched films, however, experienced
differing effects with respect to the magnitude of the wetting angle,
which increased rapidly during the first 20 min then subsequently
decreased nearly as swiftly implying that etching does not produce
surfaces similar to that of as-deposited CdTe. The rate of change in
wetting angles was greatest within the first hr of exposure to the
atmosphere for both as-deposited and etched CdTe. CdTeO~3~ is the
predominant native oxide of CdTe. Changes in the magnitude of wetting
angle can therefore be attributed to the formation of the native oxide
CdTeO~3~ for as-dep CdTe and in-part for the etched CdTe. For etched
CdTe it is hypothesized that initially a thin amorphous Te layer is
generated which begins to crystallize over a twenty-min period and then
subsequently oxidizes to TeO~2~, the predominant native oxide for Te.

The results of processing on the surface energies of PVD CdTe were
studied. The oxides that are formed or removed as a result of processing
primarily affect E^p^ leaving E^d^ largely unchanged, however, contact
wetting was unable to differentiate between the formation of different
oxide phases. Surface oxidation resulting from CdCl~2~ treatment
increased E^p^, by 9.8, 8.7, and 5.9 mN/m for vacuum, dry air, and
saturated air stored samples, respectively. Etching in Br~2~:CH~3~OH
reduced E^p^ by 4.4, 16.8, and 9.6 mN/m due to the removal of oxides
from the film surfaces for vacuum, dry air, and saturated air stored
samples, respectively.

Effects of crystallization of ITO were evaluated for atmospheres
containing air and argon. Energy changes may be attributed to higher
surface polarity associated with the alignment and organization of bonds
from previously amorphous states.

Thus, contact wetting is a simple and effective tool that is able to
detect changes in surface energies resulting from post deposition
processing. As-deposited CdTe seeks to satisfy dangling bonds through
reaction with air ambient. Significant native oxidation of PVD (111)
CdTe occurs within hours upon exposure to the atmosphere and can be
observed by increases in wetting angles. Dry air storage of CdTe prior
to CdCl~2~ heat treatments forms beneficial (CdO) oxides while films
kept in saturated air produce electronically detrimental CdTe~2~O~5~ on
surfaces. Polar energies increase as a result of oxides formed during
treatment and decrease upon removal via chemical etching.

# LIST OF CONTRIBUTORS- 9/05/2003 -- 9/04/2004

######### 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
>
> Shannon Fields
>
> **POST-DOCTORAL FELLOWS/**
>
> **LIMITED TERM RESEARCHER**
>
> Estela Calixto-Rodriguez
>
> Ujjwal Das
>
> Gautam Ganguly
>
> Greg Hanket
>
> Puthar Devassey Paulson
>
> Sergey Rykov
>
> Alex Stavrides

######### TECHNICAL SUPPORT

> Joshua Cadoret
>
> Vicky DiNetta
>
> Ronald Dozier
>
> Gautam Ganguly
>
> Kevin Hart
>
> **GRADUATE STUDENTS**
>
> Michael Burrows
>
> Darshini Desai
>
> Ozgenc Ebil
>
> Toby Junker
>
> Meijun Lu
>
> Kapil Mukati
>
> Sarah Rickman
>
> Scott Stephens
>
> Jochen Titus
>
> Clifford Yapp
>
> Mike Zellner

######### STUDENTS

> Michael Angelo
>
> Tyreef Bevel
>
> Derek Birkmire
>
> Stephanie Einstein
>
> Jonathan Frey
>
> Mary Kondos
>
> Kunal Modi
>
> Adam Muraskas
>
> Daniel Ryan
>
> David Silowka
>
> Brian Wagner

######### DOCUMENT PREPARATION

> Steven S. Hegedus
>
> L. Bonnie Johnston

# TABLE OF CONTENTS

SUMMARY [iii](#summary)

**CdTe-Based Solar Cells** [iii](#cdte-based-solar-cells)

**CuInSe~2~-Based Solar Cells** [vii](#cuinse2-based-solar-cells)

**Si-Based Solar Cells** [ix](#si-based-solar-cells)

**Diagnostics** [xi](#diagnostics)

LIST OF CONTRIBUTORS- 9/05/2003 -- 9/04/2004
[xiii](#list-of-contributors--9052003-9042004)

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

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

LIST OF TABLES [xx](#list-of-tables)

1 INTRODUCTION [1](#introduction)

**1.1 CdTe/CdSe-based Solar Cells** [1](#cdtecdse-based-solar-cells)

**1.2 CuInSe~2~-Based Solar Cells** [1](#cuinse2-based-solar-cells-1)

**1.3 Si-based Solar Cells** [2](#si-based-solar-cells-1)

**1.4 In-line Diagnostics** [3](#in-line-diagnostics)

**1.5 Training and Education** [3](#training-and-education)

**1.6 Publications** [3](#publications)

**1.7 Organization of the Report** [3](#organization-of-the-report)

2 CdTe-Based Solar Cells [4](#cdte-based-solar-cells-1)

**2.1 Improving CdS/CdTe solar cell performance**
[4](#improving-cdscdte-solar-cell-performance)

2.1.1 Junction uniformity: buffer layers and open circuit voltage
[4](#junction-uniformity-buffer-layers-and-open-circuit-voltage)

2.1.2 Effect of processing on CdTe conductivity
[9](#effect-of-processing-on-cdte-conductivity)

2.1.3 Voltage dependent photocurrent analysis
[12](#voltage-dependent-photocurrent-analysis)

2.1.4 Upper limit to V~oc~ at low temperatures
[20](#upper-limit-to-voc-at-low-temperatures)

**2.2 Contacts and Stability** [22](#contacts-and-stability)

2.2.1 Systematic investigation of the effects of humidity during stress
of CdTe devices under illumination
[22](#systematic-investigation-of-the-effects-of-humidity-during-stress-of-cdte-devices-under-illumination)

2.2.2 Transparent ZnTe:Cu contact development
[27](#transparent-zntecu-contact-development)

**2.3 High Throughput Processing** [32](#high-throughput-processing)

2.3.1 VT system modifications [32](#vt-system-modifications)

2.3.2 Impurity effects [41](#impurity-effects)

2.3.3 High throughput post-deposition processing
[44](#high-throughput-post-deposition-processing)

2.3.4 Vapor chloride treatment -- diffusion delivery
[44](#vapor-chloride-treatment-diffusion-delivery)

2.3.5 Vapor chloride treatment -- carrier gas delivery
[47](#vapor-chloride-treatment-carrier-gas-delivery)

2.3.6 Vapor chloride -- film and device results
[49](#vapor-chloride-film-and-device-results)

2.3.7 VT cell results [50](#vt-cell-results)

**2.4 CdTe Surface Chemistry** [52](#cdte-surface-chemistry)

2.4.1 The mechanism of nitric acid/phosphoric acid etching of CdTe
[52](#the-mechanism-of-nitric-acidphosphoric-acid-etching-of-cdte)

2.4.2 Understanding the chemistry of aniline etching of CdTe
[54](#understanding-the-chemistry-of-aniline-etching-of-cdte)

3 CuInSe~2~-based Solar Cells [62](#cuinse2-based-solar-cells-2)

**3.1 In-line Deposition System Improvements:**
[62](#in-line-deposition-system-improvements)

3.1.1 Cu source [62](#cu-source)

3.1.2 Se delivery [63](#se-delivery)

3.1.3 Web temperature control [65](#web-temperature-control)

3.1.4 Film cracking [65](#film-cracking)

**3.2 CuInSe~2~-based devices with increased V~oc~**
[65](#cuinse2-based-devices-with-increased-voc)

3.2.1 Cu(InAl)Se~2~ solar cells [65](#cuinalse2-solar-cells)

3.2.2 Sulfur diffusion in Cu(InGa)Se~2~
[75](#sulfur-diffusion-in-cuingase2)

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

3.3.1 Cu(InGa)Se~2~ optical characterization
[80](#cuingase2-optical-characterization)

3.3.2 Emitter layer deposition and junction characterization
[86](#emitter-layer-deposition-and-junction-characterization)

4 Si-Based Solar Cells [91](#si-based-solar-cells-2)

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

4.1.1 Post-deposition annealing experiments
[91](#post-deposition-annealing-experiments)

4.1.2 In-situ aluminum induced crystallization experiments
[93](#in-situ-aluminum-induced-crystallization-experiments)

4.1.3 Role of Al film interface [99](#role-of-al-film-interface)

**4.2 Solar Cell Development** [103](#solar-cell-development)

4.2.1 Impurities in a-Si layers [103](#impurities-in-a-si-layers)

4.2.2 Boron doped HWCVD Si films [109](#boron-doped-hwcvd-si-films)

5 In-line process diagnostics [111](#in-line-process-diagnostics)

**5.1 Introduction** [111](#introduction-1)

**5.2 Methodology** [113](#methodology)

5.2.1 Surface energy analysis [113](#surface-energy-analysis)

5.2.2 Contact wetting angle measurements
[115](#contact-wetting-angle-measurements)

5.2.3 Film preparation and characterization
[115](#film-preparation-and-characterization)

**5.3 Results and Discussion** [116](#results-and-discussion)

5.3.1 GIXRD patterns and surface energies for as-deposited and
heat-treated thin films of In~2~O~3~:SnO (ITO)
[116](#gixrd-patterns-and-surface-energies-for-as-deposited-and-heat-treated-thin-films-of-in2o3sno-ito)

5.3.2 Native oxidation of CdTe and its effect on surface energies
[117](#native-oxidation-of-cdte-and-its-effect-on-surface-energies)

5.3.3 Effects of storage environments prior to processing on CdTe GIXRD
patterns, XPS surface concentrations and surface energies.
[119](#effects-of-storage-environments-prior-to-processing-on-cdte-gixrd-patterns-xps-surface-concentrations-and-surface-energies.)

5.3.4 Effects of storage and processing on CdTe GIXRD patterns, XPS
surface concentrations, and surface energies
[120](#effects-of-storage-and-processing-on-cdte-gixrd-patterns-xps-surface-concentrations-and-surface-energies)

**5.4 Conclusion** [124](#conclusion)

6 References [125](#references)

#  {#section .unnumbered}

# LIST OF FIGURES

Figure 1 Tapping AFM (left) and interleaved EFG (right) of a working
cell. In the AFM image, the glass, SnO~2~, CdS, CdTe, and back contact
are visible from left-to-right. In the EFG image, the majority of the
potential contrast is due to CdS and CdTe, not SnO~2~ or back contact.
[6](#__RefHeading___Toc506348151)

Figure 2 X-Y-Z plot of the EFG data from Figure 1, with potential versus
position. [7](#__RefHeading___Toc506348152)

Figure 3 Line scan of potential versus position perpendicular to the
junction. [7](#__RefHeading___Toc506348153)

Figure 4 Temperature dependence of dark conductivity from 6 samples
whose treatments are given in Tables 2 or 3.
[10](#__RefHeading___Toc506348154)

Figure 5 J(V) curves for C2-53 measured at 100, 32, 10%, and dark.
[14](#__RefHeading___Toc506348155)

Figure 6 Collection function $h_{C}(V')$ for C2-53. Also shown is fit to
the 100%-10% data using Equation 6 with X~C~=21 and V~FB~=0.86V.
[14](#__RefHeading___Toc506348156)

Figure 7 J(V) curves for C2-14 measured after stress at 100, 10%, and
dark. [15](#__RefHeading___Toc506348157)

Figure 8 Collection function $h_{C}(V')$ for C2-14 after stress. Also
shown is fit to the 100%-10% data using Equation 6 with X~C~=5 and
V~FB~=0.73V. [16](#__RefHeading___Toc506348158)

Figure 9 Collection function $h_{C}(V')$ for UT966C after stress. Also
shown is fit to the 100%-10% data using Equation 6 with X~C~=6 and
V~FB~=0.81V. [17](#__RefHeading___Toc506348159)

Figure 10 Collection function $h_{C}(V')$ calculated from AMPS
simulations for different N~A~ and φ~BC~. Solid line: AMPS calculation:
dashed line: fit using Equation 6. Curves A: N~A~=1E14 cm^-3^, φ~BC~=0,
X~C~=70, V~FB~=0.85; Curves B: N~A~=1E14 cm^-3^, φ~BC~=0.4 eV, X~C~=70,
V~FB~=1.02; Curves C: N~A~=1E15 cm^-3^, φ~BC~=0.4 eV, X~C~=12,
V~FB~=0.95. Curve A was fit only up to 0.92 V, which was V~OC~.
[19](#__RefHeading___Toc506348160)

Figure 11 Electric field profile at 0.6 V bias in the 4 µm CdTe layer
from AMPS for cases of Figure 10 with φ~BC~=0.4 eV. Reverse field at the
back is due to φ~BC~. Both curves integrate to 0.36 V giving V~BI~=0.96
V. [19](#__RefHeading___Toc506348161)

Figure 12 Bias dependence of electric field at depth 0.5 and 1.0 µm from
junction from AMPS case B with N~A~=1E14 cm^-3^ and φ~BC~=0.4 eV.
[20](#__RefHeading___Toc506348162)

Figure 13a (left) and 13b (right). Temperature dependence of V~oc~ for
C2-23 (initial state) at 4 intensities. 13a shows linear fitting. 13b
shows saturation below 180K. [21](#__RefHeading___Toc506348163)

Figure 14a (left) and 14b (right) Temperature dependence of V~oc~ for
C2-12 at 4 intensities after stress at 85°C for 5 days at V~oc~ in air.
14a shows linear fitting. 14b shows saturation below 180K.
[22](#__RefHeading___Toc506348164)

Figure 15 GIXRD of IEC Au/Cu/Te-contacted devices following 500 hrs at
80°C under illumination at various humidities.
[24](#__RefHeading___Toc506348165)

Figure 16 GIXRD IEC Au/Cu/Te-contacted devices; control and following
dark and light stress at 0% humidity for 500 hrs at 80°C.
[24](#__RefHeading___Toc506348166)

Figure 17 JV curve for device with ZnTe:Cu back contact using CuSO~4~.
[28](#__RefHeading___Toc506348167)

Figure 18 Optical transmission measurements for ZnTe:Cu films for
decreasing TEA concentration. ZT30 had no TEA. ZT51 to 54 had 21, 11, 6
and 1 drop, respectively. [29](#__RefHeading___Toc506348168)

Figure 19 JV curves for device FS 20710F3-2 having 11 drops of TEA.
Measurements for light through the front (glass/SnO~2~) and back
(ITO/ZnTe:Cu) are shown. [30](#__RefHeading___Toc506348169)

Figure 20 Light JV curves for back wall (left) and front wall and dark
(right) illumination for FS 20710F3-4 before and after heat treatment at
180°C for 15 min. [30](#__RefHeading___Toc506348170)

Figure 21a (left) and 21b (right). QE curves measured at --1V, 0V, and
+0.5V bias with light bias through the glass/SnO~2~ (Figure 21a) and
through the ITO/ZnTe:Cu (Figure 21b). No TEA was used.
[31](#__RefHeading___Toc506348171)

Figure 22 Backwall QE at 0V bias for two devices with different TEA
concentrations (11 and 21 drops). [32](#__RefHeading___Toc506348172)

Figure 23 Heating and cooling profiles for source (static) and substrate
(moving) during a baseline vapor transport deposition.
[35](#__RefHeading___Toc506348173)

Figure 24 Measured (black squares) and modeled (red circles) source rate
for 780ºC, T~sou~ \< 840ºC and 20 sccm carrier gas flow rate.
[37](#__RefHeading___Toc506348174)

Figure 25 Measured accumulation rates at the substrate (red circles) and
corresponding source rates (black squares). Multiple samples at T~sou~ =
800ºC were deposited at 500ºC, 550ºC and 570ºC. All other samples were
deposited at 500ºC. [37](#__RefHeading___Toc506348175)

Figure 26 Tapping AFM images of VT112 and VT113, deposited 2.5 cm/min
and 1.25 cm/min translation speed, yielding 3 μm and 6 μm thickness,
respectively, for T~sou~ = 800ºC, T~sub~ = 500ºC, P = 20 Torr, He
carrier gas flow rate = 20 sccm, and p(O~2~) \~ 40 mTorr.
[38](#__RefHeading___Toc506348176)

Figure 27 Tapping AFM images of VT114 and VT120, deposited at T~sub~ =
500ºC and 570ºC, respectively, for T~sou~ = 800ºC, 1.25 cm/min
translation speed. [39](#__RefHeading___Toc506348177)

Figure 28 Tapping AFM images of VT121 and VT122, deposited at p(O~2~) =
40 mTorr and 80 mTorr, respectively, for T~sou~ = 800ºC, T~sub~ = 550ºC,
and 1.25 cm/min translation speed. [39](#__RefHeading___Toc506348178)

Figure 29 Optical transmission and reflection for VTD
CdTe/CdS/Ga~2~O~3~/SnO~2~/glass. [40](#__RefHeading___Toc506348179)

Figure 30 Optical absorption versus energy for VTD
CdTe/CdS/Ga~2~O~3~/SnO~2~/glass. [40](#__RefHeading___Toc506348180)

Figure 31 Light and dark J-V curves from two IEC VT samples.
[42](#__RefHeading___Toc506348181)

Figure 32 Light and dark J-V curves from an IEC VT sample and a FS
sample, which were co-processed at IEC.
[42](#__RefHeading___Toc506348182)

Figure 33 Characteristic time versus temperature for thermal (dotted)
and mass (solid) diffusion of CdCl~2~ in nitrogen for different spacing.
[47](#__RefHeading___Toc506348183)

Figure 34 Time development of CdCl~2~ partial pressure at CdTe surface
for 400°C (D~AB~ = 0.297 cm^2^/s). [47](#__RefHeading___Toc506348184)

Figure 35 CdCl~2~ partial pressure versus position though a packed bed
generator for operation at 400°C with 0.1 cm spheres.
[48](#__RefHeading___Toc506348185)

Figure 36 GIXRD patterns of VT CdTe in as-deposited and CdCl~2~ treated
conditions, showing the contrast between wet and vapor treatments.
[49](#__RefHeading___Toc506348186)

Figure 37 QE versus wavelength for samples from Table 14.
[52](#__RefHeading___Toc506348187)

Figure 38 VASE spectra of CdTe film (a) before and (b) following 1 hr
treatment in a standard aniline bath with illumination through the
substrate glass. [55](#__RefHeading___Toc506348188)

Figure 39 GIXRD, at 1° incidence, of CdTe film following 1 hr treatment
in a standard aniline bath with illumination through the substrate
glass. [56](#__RefHeading___Toc506348189)

Figure 40 GIXRD, at 1° incidence, of CdTe film following 5.5h treatment
in a standard aniline bath with illumination through the substrate
glass. [56](#__RefHeading___Toc506348190)

Figure 41 Schematic of the proposed chemical mechanism for the formation
of surface Te during aniline-based etching of CdTe under illumination.
[60](#_Ref83193172)

Figure 42 Determining the TC offset for the Cu source during a run.
[63](#__RefHeading___Toc506348192)

Figure 43 XRD (a) and GIXRD (b) spectra from selenized BSG /Mo
substrates. [68](#__RefHeading___Toc506348193)

Figure 44 XRD (a) and GIXRD (b) spectra from selenized SLG/Mo
substrates. [69](#__RefHeading___Toc506348194)

Figure 45 XRD (a) and GIXRD (b) spectra from selenized BSG /Mo/Ga
substrates. [70](#__RefHeading___Toc506348195)

Figure 46 XRD (a) and GIXRD (b) spectra from selenized SLG/Mo/Ga
substrates. [71](#__RefHeading___Toc506348196)

Figure 47 XRD (a) and GIXRD (b) spectra from selenized BSG/Mo/Cu
substrates. [72](#__RefHeading___Toc506348197)

Figure 48 XRD (a) and GIXRD (b) spectra from selenized SLG/Mo/Cu
substrates. [73](#__RefHeading___Toc506348198)

Figure 49 XRD scan of the glass/Mo substrate after a Cu(InGa)Se~2~ film
was peeled from it. The peaks are identified in Table 17.
[74](#__RefHeading___Toc506348199)

Figure 50 XRD spectra showing the (112) peak before and after reaction
in H~2~S + O~2.~ [77](#__RefHeading___Toc506348200)

Figure 51 XRD spectra showing the (112) peak before and after reaction
in H~2~S + H~2~Se + O~2.~ [78](#__RefHeading___Toc506348201)

Figure 52 AES depth profiles of Cu(InGa)Se~2~ films reacted in H~2~S,
H~2~S + O~2~, and H~2~S + H~2~Se + O~2~.
[78](#__RefHeading___Toc506348202)

Figure 53 Ga/(In+Ga) ratios from AES depth profiles of Cu(InGa)Se~2~
films reacted in H~2~S, H~2~S + O~2~, and H~2~S + H~2~Se + O~2~.
[79](#__RefHeading___Toc506348203)

Figure 54 Cu~2~Se-In~2~Se quasi-binary phase diagram taken from
Reference 49. Triangles show the compositions of Cu(InGa)Se~2~ films
used in this study. [80](#__RefHeading___Toc506348204)

Figure 55 Schematic of a coated-sphere microstructure (a) and a random
mixture of 2 phases (b). [81](#__RefHeading___Toc506348205)

Figure 56 Index of refraction spectra obtained by EMA analysis, of
Cu-In-Ga-Se alloy films with varying Cu concentration.
[83](#__RefHeading___Toc506348206)

Figure 57 Extinction coefficient spectra obtained by EMA analysis, of
Cu-In-Ga-Se alloy films with varying Cu concentration.
[83](#__RefHeading___Toc506348207)

Figure 58 Cu~2~Se-In~2~Se~3~ pseudo-binary phase diagram taken from
Reference 49. Triangles show the compositions of CuInSe~2~ films used in
this study. [84](#__RefHeading___Toc506348208)

Figure 59 Index of refraction spectra for CuInSe~2~ films obtained from
the oscillator and EMA models. [85](#__RefHeading___Toc506348209)

Figure 60 Extinction coefficient spectra for CuInSe~2~ films, obtained
from the oscillator and EMA models. [86](#__RefHeading___Toc506348210)

Figure 61 Normalized short wavelength QE used to determine thickness and
E~g~ of CdS and Cd~1‑x~Zn~x~S emitter layers deposited on Cu(InGa)Se~2~.
[88](#__RefHeading___Toc506348211)

Figure 62 XRD pattern of evaporated Si film on Al before and after
annealing. [92](#__RefHeading___Toc506348212)

Figure 63 XRD pattern of HWCVD Si film on Al before and after post
deposition annealing. [92](#__RefHeading___Toc506348213)

Figure 64 Optical micrograph through the glass substrate of Si film
deposited by HWCVD. [93](#__RefHeading___Toc506348214)

Figure 65 Optical micrograph through the glass substrate of Si film
deposited by evaporation after 90 min anneal at 500 °C.
[93](#__RefHeading___Toc506348215)

Figure 66 Nanocrystalline Si layer atop a) and between b)
polycrystalline Si islands. [94](#__RefHeading___Toc506348216)

Figure 67 Optical micrograph through the glass substrate where the dark
features are Si grains. [95](#__RefHeading___Toc506348217)

Figure 68 Raman profiles of polycrystalline Si grains and
nanocrystalline Si/Al mixture compared to that of c-Si.
[96](#__RefHeading___Toc506348218)

Figure 69 Optical micrograph of deposited film following annealing for 2
hr at 430°C. [96](#__RefHeading___Toc506348219)

Figure 70 Optical micrographs through the glass substrate of deposited
Si films on a) 100 nm, b) 250 nm, c) 500 nm and d) 1000 nm Al.
[97](#__RefHeading___Toc506348220)

Figure 71 Optical microscopy pictures of partially crystallized Si
samples taken through the glass substrate: (a) with Al layer exposed to
air for 24 hr before Si deposition, (b) with 48 hrs, and (c) with 80
days. [99](#__RefHeading___Toc506348221)

Figure 72 Ellipsometry measurements showing a progressive change in the
optical parameter ε~2~ of 100 nm thick aluminum layers even after 12
days. Inset shows the decrease in the peak over the range 1.45 to 1.7
eV. [100](#__RefHeading___Toc506348222)

Figure 73 XPS depth profile for Al2p for 100 nm thick Al film on glass.
The binding energies for Al-Al and Al-O are 73.0 eV and 75.6 eV,
respectively. Data is normalized with respect to the Al-O peak.
[101](#__RefHeading___Toc506348223)

Figure 74 XPS depth profile for O1s showing a shift to a lower energy as
the sputtering depth increases. The energies observed are in good
agreement with the data from the literature considering the data shown
here is not corrected nor normalized.
[102](#__RefHeading___Toc506348224)

Figure 75 Auger depth profiles showing the change in composition of the
aluminum oxide layers. [103](#__RefHeading___Toc506348225)

Figure 76 Nitrogen depth profiles. [104](#__RefHeading___Toc506348226)

Figure 77 Oxygen depth profiles for all samples plus P profile for 41032
(dashed line). [106](#__RefHeading___Toc506348227)

Figure 78 (a) and 78 (b). QE curves at --1V, 0V, and +0.5V bias for
devices with 20 sec (left) and 75 sec (right) p-layer.
[108](#__RefHeading___Toc506348228)

Figure 79 QE ratios for devices with 20 and 75 sec p-layers.
[108](#__RefHeading___Toc506348229)

Figure 80 Schematic of the cross-section of a CdTe solar cell detailing
thin film layers in sequential order of deposition.
[111](#__RefHeading___Toc506348230)

Figure 81 CCD image of the cross-section of a water droplet deposited
onto CdTe with left and right contact angles represented by θ~L~ and
θ~R~, respectively. [114](#__RefHeading___Toc506348231)

Figure 82 GIXRD patterns for as-deposited and treated ITO. Treatments
included air and argon anneal at 500^°^C for 10 min. The sampling depth
(1/e basis) for Cu K~ α~ X-rays in ITO at 1^o^ incidence is \~ 1000 A.
[117](#__RefHeading___Toc506348232)

Figure 83 Wetting angle of deionized water vs. time for PVD (111) CdTe
during the first hr of exposure to the atmosphere from the evacuated
deposition chamber (solid circles) and after chemical etching of room
air stored CdTe for 5 sec in a solution of Br~2~:CH~3~OH (open squares).
[118](#__RefHeading___Toc506348233)

Figure 84 GIXRD patterns for PVD CdTe stored under vacuum desiccator
(pH~2~O \< 0.1 Torr) conditions for \~2.5 months prior to processing.
Processing treatments included CdCl~2~ treatment in air at 400^o^C for
20 min, followed by etching for 5 sec in a solution of Br~2~:CH~3~OH.
The sampling depth for Cu K~ α~ X-rays in CdTe at 1^o^ incidence is \~
1160 A. [119](#__RefHeading___Toc506348234)

Figure 85 GIXRD patterns for PVD CdTe stored under desiccator (pH~2~O= 5
Torr) conditions for \~2.5 months prior to processing. Processing
treatments included CdCl~2~ treatment in air at 400^o^C for 20 min
followed by etching for 5 sec in a solution of Br~2~:CH~3~OH. The
sampling depth for Cu K~ α~ X-rays in CdTe at 1^o^ incidence is \~ 1160
A. [120](#__RefHeading___Toc506348235)

Figure 86 GIXRD patterns for PVD CdTe stored in saturated air at 25^o^C
(pH~2~O= 24 Torr) for \~1.5 months prior to processing. Processing
treatments included CdCl~2~ heat treatment in air at 400^o^C for 20 min
followed by etching for 5 sec in a solution of Br~2~:CH~3~OH. The
sampling depth for Cu K~α~ X-rays in CdTe at 1^o^ incidence is \~ 1160
A. [121](#__RefHeading___Toc506348236)

Figure 87 CCD images of cross-sections of deionized water droplets
deposited onto CdTe thin-film after CdCl~2~ treatment (a) and subsequent
etching in Br~2~:CH~3~OH (b). In (a) the drop extends beyond the field
of view to the right due to greater surface wetting.
[123](#__RefHeading___Toc506348237)

#  {#section-1 .unnumbered}

# LIST OF TABLES

Table 1 CdTe/CdS device JV parameters for VT cells with 60 nm thick HR
layers. [8](#__RefHeading___Toc506348093)

Table 2 Conductivity results for CdTe films from PVD run 41184 with no
CdCl~2~ exposure. Contacts and post deposition treatments are indicated.
[10](#__RefHeading___Toc506348094)

Table 3 Conductivity results for CdTe films from PVD run 41184 with
standard CdCl~2~ vapor treatment. Contacts and post deposition
treatments are indicated. [11](#__RefHeading___Toc506348095)

Table 4 Conductivity results for CdTe films from PVD run 41200 and 41201
with 1X and 10X CdCl~2~ vapor partial pressure, respectively. Contacts
and post deposition treatments are indicated.
[12](#__RefHeading___Toc506348096)

Table 5 Humidity conditions for each control agent at 80°C.
[23](#__RefHeading___Toc506348097)

Table 6 J-V data for Au/Cu/Te contacted IEC devices before and after
stress at 80°C for 500 hrs under illumination at various humidities.
[25](#__RefHeading___Toc506348098)

Table 7 J-V data for First Solar devices before and after stress at 80°C
for 500 hrs under illumination in various humidities.
[26](#__RefHeading___Toc506348099)

Table 8 Concentration and volume of chemicals used in galvanic solution
growth of ZnTe. [27](#__RefHeading___Toc506348100)

Table 9 Specifications of IEC vapor transport system source.
[34](#__RefHeading___Toc506348101)

Table 10 Vapor transport deposition conditions.
[43](#__RefHeading___Toc506348102)

Table 11 Saturation pressures (P~sat~) for materials in the
CdTe-CdS-CdCl~2~ system, calculated from published parametric
thermodynamic constants \[Knacke, et al, 1997\].
[46](#__RefHeading___Toc506348103)

Table 12 Thermal (*α*) and mass (*D~AB~*) diffusivity for CdCl~2~ in
nitrogen ambient. [46](#__RefHeading___Toc506348104)

Table 13 Best and average (Avg) AM 1.5 JV results for devices fabricated
from a single plate of production vapor transport deposited CdTe/CdS and
treated at 420°C using wet CdCl~2~, parallel-plate vapor CdCl~2~, packed
bed (L = 0.2 cm) vapor CdCl~2~, and no post-deposition treatment.
[50](#__RefHeading___Toc506348105)

Table 14 Best cell AM1.5 JV results for devices fabricated with CSD CdS
and VT CdTe using a single CdTe source lot (Alfa/Aesar).
[51](#__RefHeading___Toc506348106)

Table 15 Composition, by EDS, of CIGS films over 14 runs.
[64](#__RefHeading___Toc506348107)

Table 16 Orientation parameter for different peaks determined using the
Harris method with 3 (hkl) peaks. [67](#__RefHeading___Toc506348108)

Table 17 Identification of peaks in Figure 49 by comparison to JCPDS
card files for Mo, MoSe~2~, and Cu(InGa)Se~2~.
[74](#__RefHeading___Toc506348109)

Table 18 Compositions of Cu-In-Ga-Se~2~ samples for SE analysis and the
volume fraction of the α and δ~R~ phases determined by the Maxwell
Garnet and Bruggeman EMA analyses. MSE is the mean square error in the
analysis. [82](#__RefHeading___Toc506348110)

Table 19 Composition of the Cu-In-Se films used for the ellipsometry
analysis. [84](#__RefHeading___Toc506348111)

Table 20 Solar cell parameters for Cu(InGa)Se~2~ with E~g~ = 1.2 or 1.4
eV using CdS, (CdZn)S and ZnS emitter layers.
[87](#__RefHeading___Toc506348112)

Table 21 Thicknesses (d) of CdS and Cd~1‑x~Zn~x~S determined from the
normalized QE. [88](#__RefHeading___Toc506348113)

Table 22 Solar cell parameters for Cu(InGa)Se~2~ with E~g~ = 1.2 or 1.5
eV with CdS and (CdZn)S emitter layers deposited by CBD or CSD.
[89](#__RefHeading___Toc506348114)

Table 23 Effects of NH~3~ etching Cu(InGa)Se~2~ prior to CSD CdS.
[90](#__RefHeading___Toc506348115)

Table 24 Process parameters and film properties.
[98](#__RefHeading___Toc506348116)

Table 25 a-Si i-layers and devices characterized by SIMS.
[105](#__RefHeading___Toc506348117)

Table 26 J-V parameters for a-Si pin solar cells having different
p-layer times. QE is shown in Fig. 78a and 78b.
[107](#__RefHeading___Toc506348118)

Table 27 Deposition conditions and properties of boron doped HWCVD Si
films. [110](#__RefHeading___Toc506348119)

Table 28 Properties of liquids used for contact wetting analysis, where
T~b~ is the boiling temperature, χ~p~ is the polar fraction and E~S~ is
the total surface energy at room temperature.
[114](#__RefHeading___Toc506348120)

Table 29 Surface energies for as-deposited (As-Dep) and heat-treated
(HT) ITO thin films. Energies are listed as: dispersive energy (E^d^),
polar energy (E^p^) and total surface energy (E~S~). (∆E^d^) and (∆E^p^)
signify changes in dispersive and polar energies from as-deposited
conditions, respectively. [117](#__RefHeading___Toc506348121)

Table 30 XPS atomic percent surface concentrations for PVD CdTe thin
films stored under various conditions.
[122](#__RefHeading___Toc506348122)

Table 31 Surface energies for as-deposited and treated CdTe thin films
stored under various conditions. Energies are listed as: dispersive
(E^d^), polar (E^p^), total surface energy (E~S~), change in dispersive
subsequent to each treatment (∆E^d^), and change in polar subsequent to
each treatment (∆E^p^). [123](#__RefHeading___Toc506348123)

#  {#section-2 .unnumbered}

# 

# 

# 1 INTRODUCTION

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 have 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.

-   Identify 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 (September 05, 2003 to September
04, 2004) IEC provided training and education for the following: four
post-doctoral fellows; eleven graduate students; and eleven
undergraduate students. 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-Based Solar Cells

## **2.1 Improving CdS/CdTe solar cell performance**

## 

### 2.1.1 Junction uniformity: buffer layers and open circuit voltage

The working model for the main junction in CdTe/CdS solar cells assumes
a p/n heterojunction, with photocurrent (J~L~) originating in the CdTe
layer. The CdS layer is optically parasitic, thus maximum J~L~ is
achieved for minimal CdS thickness. V~oc~ is controlled by the forward
diode current, which includes terms for injected electron current,
space-charge recombination current and interface recombination current.
Device analysis of moderate-to-high efficiency CdTe/CdS cells that
considered the main junction as a single diode element in a circuit
containing a current generator, series and shunt components showed that
the diode current is dominated by recombination in the space charge
region of the absorber[^1]. External effects such as back contact
barriers, causing rollover in JV curves, were considered as additional
elements to the equivalent circuit.[^2] Due to the polycrystalline
nature of these actual devices, diode analyses can only yield
*effective* parameters. The effective diode quality factor was
determined, 1.3\<A\<2.0, indicative of recombination in the space-charge
region of the CdTe layer.^1,^[^3]^,^[^4] Extrapolation of V~oc~ versus T
to 0 K yields a built-in barrier of 1.4 eV, nearly equal to the CdTe, or
CdTe~1-x~S~x~, band gap, again consistent with the diode current being
controlled within the absorber layer. Admittance (capacitance-voltage)
analysis consistently shows low carrier density, \< 6 x 10^14^ cm^-3^,
in a wide range of CdTe devices, and drive level capacitance has shown a
high concentration of deep (E~v~ + 0.4 to E~v~ + 0.6 eV) trap
states,[^5] \~10^15^ cm^-3^.

The model and data suggest that the path for increasing V~oc~ is
reducing recombination current in CdTe: 1) increase CdTe carrier density
to shift hole quasi-Fermi level away from the mid-gap and 2) decrease
the trap density. Controlling doping in CdTe is the subject of section
2.1.2. Decreasing the trap density depends on understanding the trap
origin(s) and may be linked to deposition temperature and the action of
the CdCl~2~ treatment. Interface recombination does not appear to have a
significant effect in present-generation CdTe/CdS devices and would be
suggested by measuring diode quality factors near 1. The equivalent
circuit model, extended to parallel diodes between CdTe/CdS and
CdTe/TCO, accounts for the drop in V~oc~ as CdS thickness is reduced[^6]
but does not explain the uniqueness of CdS in forming the junction. The
diode current can be described in terms of the usual diode equation and
the fractional area of the two diode types:

$f = \frac{\text{Area}(A)}{\text{Area}(A + B)}$ **Equation 1**

$J = \left\lbrack J_{o}^{A}(1 - f) + J_{o}^{B}(f) \right\rbrack\text{exp}(\beta V) - J_{L}$
Equation 2

Solving for the open circuit voltage (*J* = 0) gives:

$V_{\text{oc}} = \frac{1}{\beta}\text{ln}\left\lbrack \frac{J_{L}}{J_{o}^{A}(1 - f) + J_{o}^{B}(f)} \right\rbrack$
Equation 3

Characterization of actual CdTe/CdS/ITO (A) and CdTe/ITO (B) cells shows
comparable diode A-factors \~ 1.7 but diode currents of $J_{o}^{A}$= 8 x
10^-8^ mA/cm^2^ and $J_{o}^{B}$= 1 x 10^-4^ mA/cm^2^, respectively. For
1% fractional area of holes in the CdS film, the *V~oc~* would be
reduced by 300 mV.

Conceptually extending this model, of parallel diodes separated by grain
boundaries, to a distribution of different quality diodes which are
electrically connected in parallel[^7] requires knowledge of the range
of the properties of individual CdTe grain junctions and the electrical
function of the grain boundaries. Furthermore, the influence of the back
contact and its preparation on the overall device behavior is needed. A
logical approach for improving V~oc~ in CdTe/CdS devices is to: 1)
identify causes and effects of non-uniform junctions at the micro-scale;
2) correlate this with gross device operation; 3) identify and utilize
deposition and post deposition processes which minimize junction
non-uniformity, thereby improving overall device performance; and 4)
evaluate alternative device elements, such as buffer layers, contact
materials and doping methods.

The cause-and-effect of localized non-uniformity can be tested by
preparation of specific device structures which allow separation of
influences, such as using a transparent back contact to bi-facially
illuminate the device and thus separate back contact from main junction
effects. Changing the CdTe thickness can be used to vary the separation
of the junctions in such diagnostic structures. This work is discussed
in section 2.2.2. The degree of spatial non-uniformity may be directly
measured using different micro-scale techniques to determine local
potential, field, or current. These measurements can be carried out with
high spatial resolution using a modified scanning probe microscope on
individual layers and completed devices in plan orientation or
cross-sections.

During this research phase, we have refined sample preparation
procedures to yield polished cross-sections of working cells and have
modified the Digital Instruments *Nanoscope*^®^ probe microscope for
operation in different electrical configurations. To prepare a device
cross-section, the cell is first fractured across its middle. The sample
is then mounted on a 45° tapered platen and the exposed surface is
mechanically polished using a sequence of grits, from 15 μm to 0.025 μm.
The polished sample is chemically polished using an extremely dilute
solution of bromine-methanol (\< 0.02 vol%). Then, electrical leads are
connected to the n- and p- sides of the cell and the entire assembly is
moved to the microscope and is located beneath the probe tip. The
tip-cantilever position is controlled by a feedback loop between the
positioning piezoelectric stack and signal from 632 nm laser light
reflected from the end of the cantilever. Electric field gradient (EFG)
measurements have been made with and without additional illumination and
at different junction bias. Figure 1 shows spatially coordinated
(interleaved) images of a working 11% VT cell in tapping mode
(morphology) and EFG mode at 0V external bias. The EFG data is
re-plotted as x-y-z in Figure 2, and a single line scan across the
junction is shown as an x-y plot in Figure 3.

![](media/image2.jpeg){width="6.3902777777777775in" height="3.90625in"}

[]{#__RefHeading___Toc506348151 .anchor}Figure 1 Tapping AFM (left) and
interleaved EFG (right) of a working cell. In the AFM image, the glass,
SnO~2~, CdS, CdTe, and back contact are visible from left-to-right. In
the EFG image, the majority of the potential contrast is due to CdS and
CdTe, not SnO~2~ or back contact.

![](media/image3.pct){width="4.6625in" height="3.5590277777777777in"}

[]{#__RefHeading___Toc506348152 .anchor}Figure 2 X-Y-Z plot of the EFG
data from Figure 1, with potential versus position.

![](media/image4.pct){width="5.375in" height="4.070138888888889in"}

[]{#__RefHeading___Toc506348153 .anchor}Figure 3 Line scan of potential
versus position perpendicular to the junction.

Measurements made at forward and reverse bias and at different
illumination conditions in this cell revealed marked changes in gradient
and have yet to be fully interpreted. When the cause of the obtained
contrast is understood, then measurements will be performed at different
areas across the junction. In addition, tests will be conducted to
evaluate the effect of etching and oxidation on the obtained contrast,
and measurements will be carried out with different etchants, to remove
surface damage and reveal grain boundaries. Additional samples have been
fabricated and are being fabricated to evaluate the effect of: 1)
stress-induced degradation; 2) buffer layer; 3) CdTe thickness; and 3)
back contact method. Selected devices are being analyzed in
plan-configuration using exposed CdTe and different processing
conditions to determine changes in the surface potential. A polished
cross-section device was fabricated and sent to the Colorado School of
Mines Physics Department for micro-scale characterization using
electroluminescence.

Incorporating a high resistance transparent (HR) layer between the CdS
and the TCO has been widely demonstrated to improve the tolerance of
V~oc~ to CdS thickness and coverage. The HR layer allows the CdTe/CdS
junction quality, with maximum V~oc~ from 800-860 mV, to be maintained
but does not *improve* V~oc~. Cells with no CdS at all, with CdTe/HR/TCO
structure have 100-200 mV higher V~oc~ than cells with CdTe/TCO
structure, but this V~oc~ is \< 700 mV. This means that with an HR
layer, CdTe/oxide junctions formed via discontinuities in the CdS film
will have lower diode current and thus, these junctions will exert less
pull-down effect on the overall V~oc~. For PVD cells, we previously
found that the choice of HR material is less important than the
thickness and sheet resistance of the layer. The V~oc~ for CdTe/CdS
cells with resistive layers of In~2~O~3~, ITO, ZnS, SnO~2~, and
Ga~2~O~3~ is always 50-100 mV or so higher than for cells with no HR
layer. The optimal HR thickness, where V~oc~ maximum is obtained with no
measurable addition to series resistance, is 50-100 nm for HR layers
formed by sputtering or native oxide conversion of sputtered metals.
During this year, baseline VT devices were fabricated with 60 nm thick
Ga~2~O~3~ HR layers on 10 x 10 cm L.O.F./Pilkington TEC-15^®^ TCO/soda
lime glass superstrates. Two runs were also carried out with 60 nm thick
In~2~O~3~ HR layers. The CdS films were deposited by chemical surface
deposition to a thickness of 100 nm. Cells were completed by CdCl~2~
vapor treatment at 410ºC for 20 min, etching and evaporation of Cu/Au
contacts. Cell data for these samples are shown in Table 1 below.

[]{#__RefHeading___Toc506348093 .anchor}Table 1 CdTe/CdS device JV
parameters for VT cells with 60 nm thick HR layers.

  ------------ ----------- ----------- ----------- ----------- -----------
  **Sample**   **HR        **V~oc~**   **J~sc~**   **FF**      **Eff**
               Layer**                                         

  VT118.1      Ga~2~O~3~   754         24.4        66.2        12.2

  VT119.2      Ga~2~O~3~   792         23.4        66.6        12.3

  VT120.3      In~2~O~3~   573         24.2        61.9        8.6

  VT121.3      In~2~O~3~   597         24.2        66.5        9.6
  ------------ ----------- ----------- ----------- ----------- -----------

> **CdTe deposition temperature, 550ºC.**

The table shows a strong V~oc~ sensitivity to the choice of HR material
for VT cells, with the other JV parameters remaining constant. For PVD
cells, with CdTe deposited at \~300ºC, equivalently high V~oc~ \> 750 mV
was obtained for HR thickness \> 50 nm, so the discrepancy here may
indicate a processing intolerance of In~2~O~3~ to the higher.

### 2.1.2 Effect of processing on CdTe conductivity 

As-deposited CdTe films do not have the required electronic properties,
specifically doping density and lifetime, needed to obtain high quality
junctions. Post deposition treatments are needed to shift the defect
equilibrium sufficiently to allow formation of acceptor complexes, which
have relatively low ionization potentials. 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.[^8] Analogous behavior may be expected in thin films. 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 in
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. We focus on temperature
dependent conductivity and the ratio of light to dark conductivity to
give insight into electronic properties of majority carriers. The effect
of contacts, high temperature anneal (HTA), CdCl~2~ treatment, and
surface treatments were studied.

The dark and light cross-grain conductivity versus T on CdTe films from
4 PVD depositions: 41184 and 41192 were typical standard depositions,
and 41200 and 41201 had CdCl~2~ vapor inside the bell-jar during the
CdTe deposition; i.e. an *in-situ* CdCl~2~ treatment. Films were 3-4 µm
thick and deposited on Corning 7059 glass. A co-planar contact mask was
used, having 0.05cm spacing. Three types of contacts were investigated
to determine if they influenced the conductivity: 50 nm evaporated Au;
30 nm Cu/50 nm Au; and 50 nm Te/30 nm Cu/50 nm Au. Conductivity was
measured as a function of temperature in the dark and then at 300K in
the light. Films were processed in three sets: the first set received no
CdCl~2~ treatments, the second set received the standard CdCl~2~
post-deposition treatment, and the third set received in-situ CdCl~2~
treatment during CdTe growth. The temperature range was 220 to 320K on
the first set and 210K to 370K on the second and third sets. N~2~
purging was used. The voltage range was selected to give measurable
currents, typically ±10, 20, 50 V.

The first series of samples received the following treatments: no
treatment; HTA ≡ Ar \@550°C for 60 min; and Te vapor \@550°C for 60 min
in a sealed ampoule at p=1E-6 Torr. No CdCl~2~ was applied to this first
group. Results are in Table 2, showing dark and light conductivity at
300K and dark conductivity activation energy. The temperature dependence
of three of these samples is shown in Figure 4. Most samples had a
slightly 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 were linear in the light. Hot probe measurements were also
made, using a hot tip at 250°C and a cold tip at room temperature.
As-deposited (untreated) 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___Toc506348094 .anchor}Table 2 Conductivity results for
CdTe films from PVD run 41184 with no CdCl~2~ exposure. Contacts and
post deposition treatments are indicated.

+--------+----------------+---------+-------+--------+-------+-------+
| **P    | **Treatment**  | **Co    | *     | *      | **Ra  | *     |
| iece** |                | ntact** | *Dark | *Light | tio** | *Ea** |
|        |                |         | co    | c      |       |       |
|        |                |         | nd.** | ond.** | **    | **(   |
|        |                |         |       |        | L/D** | eV)** |
|        |                |         | **(S/ | **(S   |       |       |
|        |                |         | cm)** | /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   | 0.57  |
|        |                | 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 |       |        |       |       |
+--------+----------------+---------+-------+--------+-------+-------+

![](media/image5.pct){width="4.0159722222222225in"
height="3.134027777777778in"}

[]{#__RefHeading___Toc506348154 .anchor}Figure 4 Temperature dependence
of dark conductivity from 6 samples whose treatments are given in Tables
2 or 3.

Three conclusions emerge from this experiment. First, the two untreated
samples (32a and 21c) have the lowest dark conductivity ( σ~D~)
independent of contact, and the highest activation energy (E~A~). This
relation is expected since σ~D~ \~ exp \[- E~A~ /kT\]. Second, both
samples treated in Te vapor 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 annealing in Te overpressure makes CdTe
more conductive. Third, it appears that annealing in Ar (HTA) increases
the light conductivity but not dark thus increasing the ratio of
σ~L~/σ~D~. There is no clear difference between samples with Au or
Te/Cu/Au contacts. Subsequent experiments used Cu/Au contacts.

The second group of samples received a standard CdCl~2~ vapor treatment
(410°C, for 20 min in Ar/O~2~). Prior to this they received treatments:
none; Ar \@550°C for 60 min (standard HTA); or Te vapor \@550°C for 60
min. The goal was to determine if samples could be made more conductive
(more p-type) with the Te treatment. Sample 21a represents the standard
device processing of 550°C HTA followed by 410°C CdCl~2~ vapor
treatment. Results are in Table 3 and figure 4.

[]{#__RefHeading___Toc506348095 .anchor}Table 3 Conductivity results for
CdTe films from PVD run 41184 with standard CdCl~2~ vapor treatment.
Contacts and post deposition treatments are indicated.

+------+--------------+----------+-----------+-----------+-----+-----+
| *    | *            | **C      | **Dark    | **Light   | **R | **E |
| *Pie | *Treatment** | ontact** | cond.**   | cond.**   | ati | a** |
| ce** |              |          |           |           | o** |     |
|      |              |          | *         | *         |     | **  |
|      |              |          | *(S/cm)** | *(S/cm)** | *   | (eV |
|      |              |          |           |           | *L/ | )** |
|      |              |          |           |           | D** |     |
+------+--------------+----------+-----------+-----------+-----+-----+
| 21b  | None         | Cu/Au    | 6E-5      | 7E-3      | 120 | 0   |
|      |              |          |           |           |     | .39 |
+------+--------------+----------+-----------+-----------+-----+-----+
| 21a  | Ar           | Cu/Au    | 2E-3      | 2E-1      | 100 | 0   |
|      | ,550°C,60min |          |           |           |     | .30 |
+------+--------------+----------+-----------+-----------+-----+-----+
| 31a  | Te           | Cu/Au    | 5E-4      | 4E-2      | 90  | 0   |
|      | ,550°C,60min |          |           |           |     | .29 |
+------+--------------+----------+-----------+-----------+-----+-----+

Comparing the data from Table 2 and Table 3, the CdCl~2~ treatment
increases the dark and light conductivity by 3-5 orders of magnitude. It
has been commonly assumed that this occurred, but the data confirms the
magnitude of the increase. However, despite the high photoconductivity,
the CdTe films are still photoconductive. Both sets of data indicate
that the HTA in Ar increases the light conductivity more than the Te
treatment; i.e. the CdTe in the light is about a factor of 10 times more
conductive with the HTA. This was unexpected and is also true for both
sets of contacts in Table 2. But a significant increase in conductivity
with little change in E~A~ implies an increase in mobility. An increase
in mobility of 3-5 orders of magnitude is huge, signifying a change in
transport mechanism. Further comparison is difficult since it is not
clear whether to compare to the Au or the Te/Cu/Au data from Table 2.

A third set of samples was deposited with *in-situ* CdCl~2~ vapor. The
concentration of CdCl~2~ was controlled via the temperature of the
CdCl~2~ source in the evaporator. Run #41201 had 500X higher CdCl~2~
partial pressure (p=11 mT) compared to run #41200 (p=2E-3 mT). The 11 mT
value is same as in our standard CdCl~2~ vapor treatment. What
differentiates films from 41201 from standard CdCl~2~ treated films is
that the CdCl~2~ treatment was done in the absence of oxygen and during
the CdTe growth. Pieces from each run were measured after no further
treatment and after HTA. The results are in table 4.

[]{#__RefHeading___Toc506348096 .anchor}Table 4 Conductivity results for
CdTe films from PVD run 41200 and 41201 with 1X and 10X CdCl~2~ vapor
partial pressure, respectively. Contacts and post deposition treatments
are indicated.

+-----+-----+--------------+-------+----------+----------+-----+----+
| **P | **  | *            | *     | **Dark   | **Light  | **R | ** |
| iec | CdC | *Post-CdTe** | *Cont | cond.**  | cond.**  | ati | Ea |
| e** | l~2 |              | act** |          |          | o** | ** |
|     | ~** | *            |       | **       | **       |     |    |
|     |     | *Treatment** |       | (S/cm)** | (S/cm)** | *   | ** |
|     |     |              |       |          |          | *L/ | (e |
|     |     |              |       |          |          | D** | V) |
|     |     |              |       |          |          |     | ** |
+-----+-----+--------------+-------+----------+----------+-----+----+
| 321 | 1X  | None         | Cu/Au | 7E-9     | 4E-7     | 57  | 0. |
|     |     |              |       |          |          |     | 63 |
+-----+-----+--------------+-------+----------+----------+-----+----+
| 322 | 1X  | Ar,550°C, 20 | Cu/Au | 1E-7     | 6E-5     | 600 | 0. |
|     |     | min          |       |          |          |     | 62 |
+-----+-----+--------------+-------+----------+----------+-----+----+
| 321 | 5   | None         | Cu/Au | 3E-10    | 2E-6     | 6   | 0. |
|     | 00X |              |       |          |          | 600 | 65 |
+-----+-----+--------------+-------+----------+----------+-----+----+
| 322 | 5   | Ar,550°C, 20 | Cu/Au | 4E-7     | 1E-4     | 250 | 0. |
|     | 00X | min          |       |          |          |     | 53 |
+-----+-----+--------------+-------+----------+----------+-----+----+

The activation energy is quite deep, similar to untreated pieces 21c and
32a from Table 2, which also had no CdCl~2~ or HTA. The light
conductivity is several orders of magnitude lower compared to pieces in
Table 3, which received a post-deposition CdCl~2~ treatment. This
suggests the *in-situ* treatment failed to activate the conductivity,
which may be due to the lack of oxygen since it has been established
that oxygen is vital effectiveness of the CdCl~2~ post-deposition vapor
treatment. Comparing data from all three tables to the HTA increases the
light conductivity by \~100X for all samples, regardless of whether or
not they received CdCl~2~ during growth. But to obtain the highest
conductivity and lowest activation energy, both CdCl~2~ and HTA are
needed (Table 3). Further experiments are planned to examine the effects
of moisture and oxygen during CdCl~2~ treatment.

### 2.1.3 Voltage dependent photocurrent analysis

It is becoming clear that standard solar cell "third level metrics",
which describe junction losses and solar cell behavior like A and J~O~
can be difficult to obtain unambiguously in CdTe, due to blocking
contacts and photoconductive effects. Further, they often have a poor
correlation with second level metrics like V~oc~, J~SC~, and FF.

Different parameters are needed to characterize device performance and
to correlate with processing. In this section and the following one, we
present two new methods to obtain parameters, which may provide a better
correlation with FF and V~oc~.

Photocurrent collection is difficult to analyze since it must be
inferred from the net current flow in the device. Electrons generated by
light are indistinguishable from those injected at forward bias.
Collection losses can be especially significant in CdTe solar cells due
to their low diffusion lengths, which leads to a relatively greater
reliance on electric field assisted drift, and the high electric field
due to the low carrier density. The electric field, and thus the
photocurrent collection efficiency, is reduced with increasing forward
bias. Thus, the photocurrent available at maximum power (J~MP~) is less
than at short circuit (J~SC~). Voltage dependent photocurrent collection
or J~L~(V) has been analyzed extensively in a-Si based solar cells.[^9]
J~L~(V) reduces the fill factor (FF) more than it reduces V~OC~ or
J~SC~. J~L~(V) losses in CdTe solar cells have been
considered.[^10]^,^[^11]^,^[^12]^,^[^13] In this work, we analyze the
J~L~(V) effect in CdTe solar cells by measuring current-voltage curves
J(V) at different intensities and using the difference to quantify
J~L~(V). The result is fit to a simple expression based on drift
collection in a *p-i-n* device with a uniform fieldError: Reference
source not found. In practice, the voltage dependent collection could be
determined from SR measurements at forward bias. However, complications
due to series resistance, contact barriers[^14] and photoconductive
CdS[^15] can make them unreliable for analysis.

The net current from a solar cell is the difference between the forward
bias diode current J~F~ and the opposing photocurrent flow J~L~, or
J(V)= J~F~ (V)- J~L~(V) which explicitly identifies both terms as
function of voltage. The photocurrent at a given intensity 1 is assumed
to be a function of voltage:

$J_{L_{1}}(V') = J_{\text{L0}_{1}}h_{C}(V')$ Equation 4

where J~LO~ is the maximum photocurrent at reverse bias giving complete
collection, and is proportional to light intensity.
$h_{C}\left( V^{'} \right)$ is the voltage dependent collection
efficiency and varies from 1 at reverse bias to 0 at forward bias. The
measured voltage V is corrected for series resistance R to give the
actual voltage across the semiconductor absorber layer for each J(V)
curve or V′=$V - \text{RJ}$. The photocurrent collection efficiency
$h_{C}\left( V^{'} \right)$ is obtained as the difference of J(V′) at
two intensities as:Error: Reference source not found

$h_{c}(V') = \left\lbrack \lbrack J_{1}(V') - J_{2}(V')\frac{\rbrack}{\lbrack}J_{\text{LO}1} - J_{\text{LO}2}\rbrack \right\rbrack$
. Equation 5

assuming J~F~(V) is independent of intensity.

CdTe/CdS devices from First Solar, University of Toledo, Colorado State
University, and Matsushita were analyzed. Measurements were performed
with a Xe solar simulator using neutral density filters to reduce the
intensity as needed. Figure 5 shows the J(V) curves for a CdS/CdTe
device C2-53 measured at 4 intensities (100%, 32%, 10% and dark). At
100% light, the photovoltaic parameters were V~oc~=0.82V, J~sc~=19.3
mA/cm^2^, FF=68% and efficiency=10.7%. Note the increasing curvature and
decreasing slope beyond V~oc~ with decreasing intensity. This light to
dark crossover, or LDXO, is commonly seen in CdTe devices and has been
attributed to photoconductive CdS.Error: Reference source not found The
J(V) data was corrected for R[^16] and the J(V') data at 32, 10% and
dark subtracted from the 100% data, normalized according to equation 5.
The result is plotted in figure 6. The different intensities all give
the same voltage dependent collection function $h_{C}(V')$. This
verifies the assumption that J~F~(V) and $h_{C}(V')$ are independent of
intensity at least up to where $h_{C}(V')$=0. At maximum power (MP), the
photocurrent has been reduced by about 8% representing a similar loss in
FF.

![](media/image6.pct){width="3.245138888888889in" height="3.15in"}

[]{#__RefHeading___Toc506348155 .anchor}Figure 5 J(V) curves for C2-53
measured at 100, 32, 10%, and dark.

![](media/image7.pct){width="3.3722222222222222in"
height="3.1506944444444445in"}

[]{#__RefHeading___Toc506348156 .anchor}Figure 6 Collection function
$h_{C}(V')$ for C2-53. Also shown is fit to the 100%-10% data using
Equation 6 with X~C~=21 and V~FB~=0.86V.

It is well established that some CdTe devices degrade when exposed to
accelerated stress testing, [^17]^,^[^18] which has been attributed to
increased recombination, resistance, and contact barriers.[^19] A device
C2-14 cut from the same module and having comparable initial performance
to C2-53 was stressed at 85°C for 30 days under 1 sun illumination at
--0.5V. At 100% light, the photovoltaic parameters after stress were
V~oc~=0.77V, J~sc~=17.7 mA/cm^2^, FF=53% and efficiency=7.3%. Compared
to C2-53, the largest decrease was in FF (22% relative). Figure 7 shows
the J(V) data at three intensities. An increase in LDXO and forward bias
curvature is apparent as commonly seen after stress, $h_{C}(V')$ is
shown in figure 8. Once again, the data reduces to a single voltage
dependent collection function. It was fit only up to the voltage where
significant curvature due to the blocking contact distorts the
collection function. At MP, the photocurrent has then been reduced by
22%, an 18% relative decrease indicating the decrease in FF after stress
is largely due to a decrease in $h_{C}(V')$. Thus, the loss in FF due to
decreased $h_{C}(V')$ is substantial.

![](media/image8.pct){width="3.245138888888889in" height="3.15in"}

[]{#__RefHeading___Toc506348157 .anchor}Figure 7 J(V) curves for C2-14
measured after stress at 100, 10%, and dark.

![](media/image9.pct){width="3.15in" height="2.998611111111111in"}

[]{#__RefHeading___Toc506348158 .anchor}Figure 8 Collection function
$h_{C}(V')$ for C2-14 after stress. Also shown is fit to the 100%-10%
data using Equation 6 with X~C~=5 and V~FB~=0.73V.

Calculations of voltage dependent collection losses in CdTe solar cells
using a standard model of a *p-n* junction depletion width and bulk
diffusion found a \~2% reduction in J~L~(V) at V~OC~[^20] due to the
voltage dependence of the depletion width. Figures 6 and 7 indicate a
loss of \>50% at V~OC~. Clearly other mechanisms are responsible.
Analysis of field-aided drift collection in a-Si *p-i-n* solar cells has
shown that $h_{C}(V')$ data could be fit with the expression:Error:
Reference source not found

$h_{c}(V') = X(V')\left\lbrack 1 - \text{exp}\left( X(V')^{- 1} \right) \right\rbrack$
. Equation 6

Under the assumption of a uniform electric field, it was shown that the
parameter X(V') can be defined as:[^21]

$X(V') = X_{C}\left( 1 - \frac{V'}{V_{\text{FB}}} \right)$ Equation 7

where X~C~ is the ratio of zero bias drift length to thickness and V~FB~
is the flat band voltage making field F(V) go to zero when applied or:

F(V)=k(V- V~FB~) Equation 8

where k is a constant.

This model, contained in Equations 6-8, was widely and successfully
applied to characterize voltage dependent collection in a-Si p-i-n solar
cells and modules. Equation 6 was fit to the collection function data of
Figures 6 and 7 using X~C~ and V~FB~ as adjustable parameters. V~FB~ is
determined by the voltage axis intercept and has little freedom for
adjustment, thus X~C~ is the primary fitting parameter. As X~C~
decreases, J~L~(V) losses increase,$h_{C}(V')$ becomes more rounded, and
FF decreases.Error: Reference source not found Very acceptable agreement
is shown between the analytical expression and data. The fitting
parameters are given in the caption. The decrease in FF from 68 to 53%
corresponding to a decrease in X~C~ from 21 to 5 is in excellent
quantitative agreement with a-Si devicesError: Reference source not
found suggesting that the FF in CdTe has a similar dependence on J~L~(V)
effects. This is reasonable since CdTe is nearly intrinsic like a-Si.
For completeness, Figure 9 shows $h_{C}(V')$ for a sputtered device
UT966C made at University of Toledo after stressing for 30 days at 60°C.
The fit to the data is excellent with X~C~=6. This confirms that the
approach developed here is applicable to different types of devices.

![](media/image10.pct){width="3.245138888888889in"
height="3.1479166666666667in"}

[]{#__RefHeading___Toc506348159 .anchor}Figure 9 Collection function
$h_{C}(V')$ for UT966C after stress. Also shown is fit to the 100%-10%
data using Equation 6 with X~C~=6 and V~FB~=0.81V.

Further insight into $h_{C}(V')$was obtained using numerical
simulations. J(V) curves were modeled with the AMPS analysis package
using input parameters similar to those previously established.Error:
Reference source not found^,^[^22] J(V) curves were simulated as a
function of the density of acceptor-like defects N~A~ at 0.32 eV from
the valence band and the back contact barrier φ~BC~. Figure 10 shows
$h_{C}(V')$ for three pairs of N~A~ and φ~BC~ along with the fit using
Equation 6. As N~A~ increased (compare curve B to C) with φ~BC~=0.4 eV,
the FF decreased from 81.7 to 74.5%, X~C~ decreased from 70 to 12 and
the photocurrent loss due to J~L~(V) increased from 10 to 16% at MP. The
intercept V~FB~ increases with N~A~ because the increased charge density
screens the front junction field from band bending due to φ~BC~.
Comparing curves A to B shows that increasing φ~BC~ reduces V~FB~ hence
the effective field for collection without affecting the collection
parameter since X~C~=70 in both cases. These simulations confirm that
J~L~(V) effects are not due to series resistance (there was none) or
back barriers creating an artificial light intensity dependence
resulting in an apparent J~L~(V) effect.

Earlier work Error: Reference source not
found^,Error:\ Reference\ source\ not\ found,Error:\ Reference\ source\ not\ found^
assumed the voltage dependent collection in CdTe was due to interface
recombination. We consider this unlikely for two reasons. First, we
obtained a very good fit to $h_{C}(V')$ applying the uniform field model
based on bulk collection in the absorber and second, AMPS does not
explicitly have interface states yet gave the same functional dependence
as found experimentally.

It is surprising that the uniform field model fits so well since AMPS
indicates very non-uniform field distribution. Figure 11 shows the
electric field in the CdTe layer from AMPS for the two cases of N~A~=
1E14 and 1E15 cm^-3^ (curves B and C of Figure 10) with φ~BC~=0.4 eV.
The integral of F(V,x) across the CdTe layer gave essentially V~FB~ -V;
i.e. it gave 0.96 V at 0 V bias and 0.36 V at 0.6 V bias. Thus, V~FB~ is
nearly the same as the built-in potential V~BI~ confirming one of the
basic premises of the field aided collection. The only exception was for
low N~A~ and large φ~BC~ when the front junction field extended to the
back (no screening) and was influenced by φ~BC~, leading to V~FB~ \<
V~BI~. As discussed elsewhereError: Reference source not found it is
likely the average field that governs the collection and the assumption
of uniformity is required only for simplifying the analytical
expression.Error: Reference source not found Figure 12 shows the field
at 0.5 and 1.0 µm into the CdTe. These dimensions were selected since
\>90% of the solar spectrum is absorbed within 1 µm so it is the field
in that region that is of interest to characterize the minority carrier
(electron) transport as they travel to the CdS. Figure 12 shows that the
field is nearly linear with voltage, and varies by less than 50% over
the range of voltages of interest within the absorption region. These
also support the approximation of a uniform field.

![](media/image11.pct){width="3.011111111111111in" height="2.75625in"}

[]{#__RefHeading___Toc506348160 .anchor}Figure 10 Collection function
$h_{C}(V')$ calculated from AMPS simulations for different N~A~ and
φ~BC~. Solid line: AMPS calculation: dashed line: fit using Equation 6.
Curves A: N~A~=1E14 cm^-3^, φ~BC~=0, X~C~=70, V~FB~=0.85; Curves B:
N~A~=1E14 cm^-3^, φ~BC~=0.4 eV, X~C~=70, V~FB~=1.02; Curves C: N~A~=1E15
cm^-3^, φ~BC~=0.4 eV, X~C~=12, V~FB~=0.95. Curve A was fit only up to
0.92 V, which was V~OC~.

![](media/image12.pct){width="3.8555555555555556in" height="3.15in"}

[]{#__RefHeading___Toc506348161 .anchor}Figure 11 Electric field profile
at 0.6 V bias in the 4 µm CdTe layer from AMPS for cases of Figure 10
with φ~BC~=0.4 eV. Reverse field at the back is due to φ~BC~. Both
curves integrate to 0.36 V giving V~BI~=0.96 V.

![](media/image13.pct){width="3.4444444444444446in"
height="3.1506944444444445in"}

[]{#__RefHeading___Toc506348162 .anchor}Figure 12 Bias dependence of
electric field at depth 0.5 and 1.0 µm from junction from AMPS case B
with N~A~=1E14 cm^-3^ and φ~BC~=0.4 eV.

We have analyzed over 10 CdTe devices from 4 different sources having FF
from 50 to 72 and thicknesses from 1.8 to 5 µm. All show J~L~(V) effects
similar to figures 6,8, and 9 and can be fit with equation 6. Thus, the
concept of voltage dependent collection and the *p-i-n* model appears
universally applicable to CdTe solar cells. There are questions
regarding assumption of a uniform field, neglect of diffusion, whether
X~C~ represents hole or electron transport and the role of the CdTe
thickness. This will be the subject of future effort.

In summary, voltage dependent collection in CdTe solar cells represents
a significant loss in FF, which increases after stressing. This effect
is well described using the uniform field model developed for *p-i-n*
solar cells. A single parameter X~C~ can be used to characterize the
magnitude of the loss. Numerical simulations support the experimental
results and indicate a correlation between FF, X~C~ and defects, and
between V~FB~ and built-in potential.

### 2.1.4 Upper limit to V~oc~ at low temperatures 

J(V) curves of most thin film solar cells become distorted at low
temperatures due to blocking contacts and increased resistance as is
well known. It becomes impossible to fit the J(V) curve with a single
diode model because of voltage dropped across these parasitic elements.
However, at open circuit, J=0 so there is no effect of parasitic diodes
or resistors.Error: Reference source not found In that case, V~oc~ is
well known to be described by:

V~OC~ = φ~Β~/q -- \[AkT/q\] *ln* (J~00~/J~L~) Equation 9

where J~00~ is the recombination current density prefactor, often given
as J~00~= qpS where p is the hole density and S is the recombination
velocity. Equation 9 confirms the well-established linear dependence of
V~oc~ on T, which has been shown to apply to all thin film
heterojunction solar cells. In particular, the intercept at T= 0K is the
bandgap of the absorber.Error: Reference source not found The slope
depends on the intensity and the recombination mechanism and is
typically negative 1-3 mV/K.

![](media/image14.pct){width="2.8055555555555554in"
height="2.5791666666666666in"}
![](media/image15.pct){width="2.7979166666666666in"
height="2.5972222222222223in"}

[]{#__RefHeading___Toc506348163 .anchor}Figure 13a (left) and 13b
(right). Temperature dependence of V~oc~ for C2-23 (initial state) at 4
intensities. 13a shows linear fitting. 13b shows saturation below 180K.

Figures 13a and 13b show V~oc~ (T) for the unstressed CdTe device C2-23,
which was from the same plate and had the same processing as C2-53
analyzed in the previous section. Results are shown at 4 intensities
from 6% to 100% from 220 to 315 K. Figure 13a shows linear fit from 220
to 320K. The intercepts are \~1.5 eV consistent with the CdTe bandgap.
The slopes increase with intensity as predicted by Equation 9. Figure
13b shows data down to 150K. At about 180K, the V~oc~ becomes
independent of light intensity and temperature, saturating at about
V~oc~=1.01 V. We suggest that this upper limit represents a fundamental
parameter, the built-in voltage. Note that 1 V is very close to the
value reported in the previous section from integrating the electric
field across the CdTe for typical AMPS parameters. Saturation of V~oc~
indicates carrier freeze-out. When there are no longer sufficient free
carriers, the quasi-Fermi levels no longer respond to changes in carrier
generation or temperature. V~oc~ reaches its maximum value determined by
the absorber bandgap and the contact potentials.

![](media/image16.pct){width="2.689583333333333in"
height="2.4965277777777777in"}
![](media/image17.pct){width="2.689583333333333in"
height="2.4965277777777777in"}

[]{#__RefHeading___Toc506348164 .anchor}Figure 14a (left) and 14b
(right) Temperature dependence of V~oc~ for C2-12 at 4 intensities after
stress at 85°C for 5 days at V~oc~ in air. 14a shows linear fitting. 14b
shows saturation below 180K.

Figures 14a and 14b show the same data except for a device CS-12
stressed 85°C for 5 days under 1 sun illumination at V~oc~. The V~oc~ at
300K decreased from 0.82 to 0.78V. Comparing figures 10a and 10a, the
slopes and intercepts become more light-dependant after stress. The
intercept at 1 sun intensity deceases from 1.52 to 1.43 V, consistent
with what we have observed numerous times following stress. Figure 14b
shows that the saturation value decreased by nearly the same amount,
from 1.01 to 0.93 V. This suggests the decrease in V~oc~ with stress is
not only due to greater recombination but also a change in the band
alignments or contact barriers. AMPS does not provide an adequate
representation of the J(V) curves at such low temperatures hence is not
useful to understand these effects. Further experimental work on low
temperature V~oc~ measurement of other CdS/CdTe device structures such
as having ZnTe:Cu or other back contacts and perhaps different window
layers should give insight into the meaning of the saturated value of
V~oc~.

## **2.2 Contacts and Stability** 

### 2.2.1 Systematic investigation of the effects of humidity during stress of CdTe devices under illumination

The systematic study of humidity during stress of CdTe devices and
device components has continued with the investigation of the effect of
illumination during stress at varying humidities. A new set of samples
was selected; IEC devices with 30 nm Cu/50 nm Au contacts on BDH
(Br~2~:dichrol:hydrazine) etched, which generates a ∼50nm Te surface
layer, VTD CdTe (Au/Cu/Te), and First Solar (FS) "dot" devices. Humidity
conditions of ∼0%, ∼10%, ∼50%, ∼80% and 100% were selected. Experimental
details were reported in the previous annual report. The devices were
stored in sealed quartz tubes over saturated salt solutions of LiCl,
NaBr and KBr, to produce 10%, 50% and 80% relative humidity conditions,
respectively, over Drierite pieces for the ∼0% humidity atmosphere, and
over pure water for the 100% atmosphere. All vessels were stored in an
oven at 80°C for 500 hrs over an ELH bulb, with sample vessels arranged
to ensure equal illumination, ∼0.5 suns, for all devices. Table 5
presents the salts and relative humidity conditions obtained at 80°C,
including saturation pressure and concentration of H~2~O vapor in the
vessel atmosphere.

[]{#__RefHeading___Toc506348097 .anchor}Table 5 Humidity conditions for
each control agent at 80°C.

No visual changes were observed for most pieces following illuminated
stress, with the exception of the 100% humidity FS device, which showed
some discoloration of the CdTe, and the 0% humidity IEC device, which
exhibited discoloration of the CdTe and of the back contacts, which had
become green in color.

Figure 15 shows the GIXRD patterns of the IEC devices. The formation of
TeO~2~ is observed at 0% humidity. No TeO~2~ phases are observed at
higher humidities, but it is likely that any oxides formed may be
amorphous hydroxide or hydrated phases. The diffraction pattern for the
Cu~2-x~Te back contact is observed at all humidities, suggesting that if
any oxidation does occur at the higher humidities, this phase is not
significantly degraded, though some recrystallization or re-orientation
may have occurred. The formation of a new phase, however, is observed in
the pattern for the 0% humidity stressed device. Figure 16 shows the
GIXRD plots of IEC devices; control, and stressed at 0% humidity in dark
and light. The new peaks observed in the plot of the device stressed
under illumination are assigned to a copper (ff) tellurate, CuTeO~3~
phase. The formation of CuTeO~3~ will account for the green color of the
back contacts following stress. From Figure 16, it is clear that the
formation of CuTeO~3~ at 0% humidity under illumination occurs at the
expense of the Cu~2-x~Te phase. To our knowledge, this is the first
report of the formation of CuTeO~3~ in the back contacts of stressed
CdTe devices.

##### ![](media/image18.png){width="4.054861111111111in" height="2.95625in"}

[]{#__RefHeading___Toc506348165 .anchor}Figure 15 GIXRD of IEC
Au/Cu/Te-contacted devices following 500 hrs at 80°C under illumination
at various humidities.

![](media/image19.png){width="5.340972222222222in"
height="2.7534722222222223in"}

[]{#__RefHeading___Toc506348166 .anchor}Figure 16 GIXRD IEC
Au/Cu/Te-contacted devices; control and following dark and light stress
at 0% humidity for 500 hrs at 80°C.

J-V analysis showed the IEC devices to be quite stable, despite the
harsh stress conditions. Table 6 lists the J-V data for the IEC devices
before and after illuminated humidity stress. A number of the IEC cells
exhibited forward-reverse retrace hysteresis prior to stressing, which
became worse with stressing. The other major change occurred for the
dark curve of the 100% humidity stressed device, which exhibited a high
degree of shorting after stress, but this was not mirrored in the light
trace. Overall, some decrease in V~oc~, 10-60 mV, was observed for the
devices stressed at humidities \> 0%, with greater losses at higher
humidities. J~sc~ and R~oc~ were stable and no forward bias rollover was
observed for any stressed IEC device. Interestingly, the 0% humidity
stressed device showed the best stability despite the significant
chemistry occurring in back contact during stressing.

[]{#__RefHeading___Toc506348098 .anchor}Table 6 J-V data for Au/Cu/Te
contacted IEC devices before and after stress at 80°C for 500 hrs under
illumination at various humidities.

<table>
<colgroup>
<col style="width: 10%" />
<col style="width: 10%" />
<col style="width: 5%" />
<col style="width: 7%" />
<col style="width: 7%" />
<col style="width: 7%" />
<col style="width: 6%" />
<col style="width: 6%" />
<col style="width: 4%" />
<col style="width: 4%" />
<col style="width: 6%" />
<col style="width: 7%" />
<col style="width: 7%" />
<col style="width: 7%" />
</colgroup>
<tbody>
<tr class="odd">
<td><blockquote>
<p><strong>Sample</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>condition</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>T °C</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>RH%</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>V<sub>oc</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>J<sub>sc</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>FF</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Eff</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>R<sub>oc</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>G<sub>sc</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Light retrace</strong></p>
<p><strong>Hyst.</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Dark Retrace</strong></p>
<p><strong>Hyst</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Light / dark</strong></p>
<p><strong>XOVR</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Fwd Bias Curvature</strong></p>
</blockquote></td>
</tr>
<tr class="even">
<td>VT75.43</td>
<td>Init</td>
<td></td>
<td></td>
<td>687</td>
<td>20.0</td>
<td>60.6</td>
<td>8.3</td>
<td>5</td>
<td>3</td>
<td>N</td>
<td>Y</td>
<td>N</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT75.43</td>
<td>Witness</td>
<td>25</td>
<td>30</td>
<td>707</td>
<td>20.8</td>
<td>59.9</td>
<td>8.8</td>
<td>6</td>
<td>2</td>
<td>N</td>
<td>Y</td>
<td>N</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT75.33</td>
<td>Init</td>
<td></td>
<td></td>
<td>707</td>
<td>22.1</td>
<td>61.7</td>
<td>9.6</td>
<td>5</td>
<td>2</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT75.33</td>
<td>Drierite</td>
<td>80</td>
<td>0</td>
<td>721</td>
<td>22.2</td>
<td>55.5</td>
<td>8.9</td>
<td>6</td>
<td>5</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT75.13</td>
<td>Init</td>
<td></td>
<td></td>
<td>711</td>
<td>20.9</td>
<td>61.5</td>
<td>9.1</td>
<td>7</td>
<td>2</td>
<td>N</td>
<td>W</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT75.13</td>
<td>LiCl</td>
<td>80</td>
<td>10.5</td>
<td>700</td>
<td>21.2</td>
<td>56.5</td>
<td>8.4</td>
<td>7</td>
<td>3</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT75.12</td>
<td>Init</td>
<td></td>
<td></td>
<td>704</td>
<td>21.8</td>
<td>59.5</td>
<td>9.1</td>
<td>8</td>
<td>3</td>
<td>N</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT75.12</td>
<td>NaBr</td>
<td>80</td>
<td>51.4</td>
<td>641</td>
<td>21.9</td>
<td>57.3</td>
<td>8.0</td>
<td>6</td>
<td>5</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT75.21</td>
<td>Init</td>
<td></td>
<td></td>
<td>660</td>
<td>22.7</td>
<td>51.3</td>
<td>7.7</td>
<td>7</td>
<td>6</td>
<td>Y</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT75.21</td>
<td>KBr</td>
<td>80</td>
<td>79</td>
<td>605</td>
<td>22.9</td>
<td>53.8</td>
<td>7.5</td>
<td>4</td>
<td>5</td>
<td>Y</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>VT75.32</td>
<td>Init</td>
<td></td>
<td></td>
<td>687</td>
<td>20.6</td>
<td>59.5</td>
<td>8.4</td>
<td>7</td>
<td>2</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>VT75.32</td>
<td>H<sub>2</sub>O</td>
<td>80</td>
<td>100</td>
<td>654</td>
<td>19.8</td>
<td>47.2</td>
<td>6.1</td>
<td>8</td>
<td>4</td>
<td>Y</td>
<td>Y</td>
<td>Y</td>
<td>N</td>
</tr>
</tbody>
</table>

No back contact changes were detected by GIXRD for the FS stressed
devices. Table 7 presents the J-V data for the FS device before and
after illuminated humidity stress. All stressed FS devices exhibit J-V
rollover at forward bias, increasing in severity with higher stress
humidity. Losses in V~oc~, 20-80 mV, and J~sc~, 2-6 mA, both increasing
with greater humidity, were also observed.

The chemistry of the devices, except for the formation of CuTeO~3~ at 0%
humidity under illumination, is very similar to that observed for the
pieces stressed in the dark. Likewise, comparing the J-V data of the
dark and illuminated humid stressed devices shows similar trends between
the two stress conditions, though more rapid degradation is observed for
the devices stressed under illumination.

These results highlight differences in stability of different back
contacts. In the previous reports we have attributed the thicker Te
component of the IEC back contact to the improved stability of these
devices. The thicker Te may maintain an improved contact, as it is
itself a low resistance material. Te can getter diffusing Cu ions and,
therefore, a thicker Te component may maintain a copper telluride
contact for a longer period. A third possibility is that the excess Te
may itself oxidize quicker in the harsh stress conditions and protect
the copper telluride back contact. Further chemical analysis of the
devices is required to confirm if oxidation of the back contacts does
occur at high humidity conditions. Extended stressing will be carried
out on IEC devices to observe the long-term effects of CuTeO~3~
formation on back contact performance and stability.

[]{#__RefHeading___Toc506348099 .anchor}Table 7 J-V data for First Solar
devices before and after stress at 80°C for 500 hrs under illumination
in various humidities.

<table>
<colgroup>
<col style="width: 8%" />
<col style="width: 10%" />
<col style="width: 5%" />
<col style="width: 6%" />
<col style="width: 5%" />
<col style="width: 6%" />
<col style="width: 6%" />
<col style="width: 6%" />
<col style="width: 5%" />
<col style="width: 5%" />
<col style="width: 8%" />
<col style="width: 8%" />
<col style="width: 8%" />
<col style="width: 8%" />
</colgroup>
<tbody>
<tr class="odd">
<td><blockquote>
<p><strong>Sample</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Conditions</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>T°C</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>RH%</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>V<sub>OC</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>J<sub>SC</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>FF</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Eff</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>R<sub>OC</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>G<sub>SC</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Dark Retrace</strong></p>
<p><strong>Hyst.</strong></p>
</blockquote></td>
<td><blockquote>
<p>Light Retrace</p>
<p><strong>Hyst.</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Light /Dark</strong></p>
<p><strong>XOVR</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Fwd bias</strong></p>
<p><strong>Curvature</strong></p>
</blockquote></td>
</tr>
<tr class="even">
<td>C2.43</td>
<td>Init</td>
<td></td>
<td></td>
<td>830</td>
<td>19.2</td>
<td>66.7</td>
<td>10.6</td>
<td>9</td>
<td>1</td>
<td>W</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>C2.43</td>
<td>Witness</td>
<td>25</td>
<td>30</td>
<td>818</td>
<td>18.1</td>
<td>64.2</td>
<td>9.5</td>
<td>10</td>
<td>1</td>
<td>N</td>
<td>N</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="even">
<td>C2.41</td>
<td>Init</td>
<td></td>
<td></td>
<td>829</td>
<td>19.3</td>
<td>66.5</td>
<td>10.7</td>
<td>8</td>
<td>1</td>
<td>W</td>
<td>W</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>C2.41</td>
<td>Drierite</td>
<td>85</td>
<td>0</td>
<td>808</td>
<td>18.7</td>
<td>59.1</td>
<td>8.9</td>
<td>19</td>
<td>2</td>
<td>W</td>
<td>W</td>
<td>Y</td>
<td>Y</td>
</tr>
<tr class="even">
<td>C2.42</td>
<td>Init</td>
<td></td>
<td></td>
<td>827</td>
<td>16.5</td>
<td>65.9</td>
<td>9.0</td>
<td>10</td>
<td>1</td>
<td>W</td>
<td>W</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>C2.42</td>
<td>LiCl</td>
<td>85</td>
<td>10.5</td>
<td>789</td>
<td>18.4</td>
<td>53.1</td>
<td>7.7</td>
<td>32</td>
<td>0</td>
<td>W</td>
<td>W</td>
<td>Y</td>
<td>Y</td>
</tr>
<tr class="even">
<td>C2.54</td>
<td>Init</td>
<td></td>
<td></td>
<td>832</td>
<td>19.4</td>
<td>64.6</td>
<td>10.4</td>
<td>9</td>
<td>0</td>
<td>W</td>
<td>W</td>
<td>Y</td>
<td>N</td>
</tr>
<tr class="odd">
<td>C2.54</td>
<td>NaBr</td>
<td>85</td>
<td>51.4</td>
<td>772</td>
<td>17.9</td>
<td>43.9</td>
<td>6.1</td>
<td>103</td>
<td>3</td>
<td>N</td>
<td>W</td>
<td>Y</td>
<td>Y</td>
</tr>
<tr class="even">
<td>C2.53</td>
<td>Init</td>
<td></td>
<td></td>
<td>816</td>
<td>19.3</td>
<td>67.0</td>
<td>10.6</td>
<td>7</td>
<td>1</td>
<td>N</td>
<td>N</td>
<td>W</td>
<td>N</td>
</tr>
<tr class="odd">
<td>C2.53</td>
<td>KBr</td>
<td>85</td>
<td>79</td>
<td>782</td>
<td>17.1</td>
<td>20.4</td>
<td>2.7</td>
<td>654</td>
<td>11</td>
<td>N</td>
<td>W</td>
<td>Y</td>
<td>Y</td>
</tr>
<tr class="even">
<td>C2.52</td>
<td>Init</td>
<td></td>
<td></td>
<td>816</td>
<td>19.3</td>
<td>67.2</td>
<td>10.6</td>
<td>7</td>
<td>1</td>
<td>W</td>
<td>W</td>
<td>W</td>
<td>N</td>
</tr>
<tr class="odd">
<td>C2.52</td>
<td>H<sub>2</sub>O</td>
<td>85</td>
<td>100</td>
<td>735</td>
<td>12.9</td>
<td>19.2</td>
<td>1.8</td>
<td>542</td>
<td>18</td>
<td>N</td>
<td>W</td>
<td>Y</td>
<td>Y</td>
</tr>
</tbody>
</table>

### 2.2.2 Transparent ZnTe:Cu contact development

Development of ZnTe:Cu contacts at NREL and elsewhere has been motivated
by the high p-type conductivity of the ZnTe, needed for low contact
resistance, and as a source of Cu doping the CdTe, needed for high
V~oc~. In addition, we are interested in this material as a
semi-transparent back contact for two additional reasons. First,
bifacial JV and QE measurements will provide new insight on the carrier
transport, device physics and back junction properties. Second, a
transparent contact is needed for the top cell of a polycrystalline
tandem based on CdTe.

Previously, IEC had developed galvanic solution growth of ZnTe:Cu films
for use as a low resistance back contact and source of Cu doping.[^23]
During the present contract period, IEC investigated the solution growth
of ZnTe based on our previous work, pursuing two chemical pathways for
the source of Cu: CuCl~2~ and CuSO~4~. Devices with CuSO~4~ achieved
consistently higher efficiency and hence further work focused on using
CuSO~4~ as the Cu source using CdS/CdTe substrates from First Solar and
from IEC's vapor transport system. Initial development used substrates
from First Solar modules, which received the wet CdCl~2~ process and BDH
surface treatment at IEC. After a 10-min dip in the ZnTe solution to
form the primary ZnTe:Cu contact, cells received ITO and Ni/Al grids to
form the transparent secondary back contact. Variations in solution
concentration and time were evaluated by measuring the absorption and
conductivity and by making contacts to solar cells. Based on those
results, the following chemical constituents and concentrations were
selected for further optimization:

The pH was maintained at 3 and temperature during deposition was 68°C.

[]{#__RefHeading___Toc506348100 .anchor}Table 8 Concentration and volume
of chemicals used in galvanic solution growth of ZnTe.

  ----------------------- ------------------------- ---------------------
  **Solution**            **Concentration (M)**     **Volume (mL)**

  ZnSO~4/~ZnCl~2~         0.1M                      54

  CuSO~4/~ZnCl~2~         0.001M                    18

  TeO~2~                  0.001M                    16
  ----------------------- ------------------------- ---------------------

Devices with CuSO~4~ achieved 10% efficiency, but had hysteresis in the
J-V retrace and slight shunting and rollover in forward bias, suggesting
too much free Cu and/or formation of CuTe.

![](media/image20.pct){width="3.25in" height="3.1590277777777778in"}

[]{#__RefHeading___Toc506348167 .anchor}Figure 17 JV curve for device
with ZnTe:Cu back contact using CuSO~4~.

Samples were given heat treatments totaling 30 min at 180°C in Ar, which
decreased the hysteresis and increased V~oc~ by 0.1-0.2 V. The light and
dark JV curve for the best cell with the standard CuSO~4~ solution after
heat treatment is shown in Figure 17. The parameters for the up-trace
(higher maximum power) were V~oc~=0.755V, J~sc~=21.4 mA/cm^2^, FF=62.7%,
and efficiency=10.1%. The down-trace had FF=58%. The dark JV curve had
more severe instability than the light curve, as is often observed in
both CdTe and Cu (InGa) Se~2~. Cells with CdCl~2~ achieved only about
7.5%, having lower FF and even more hysteresis than cells with CuSO~4~
as the primary source of Cu. The back wall J~sc~, through the
ITO/ZnTe:Cu back contact, was very low, \<1 mA/cm^2^. Optical and QE
measurements indicated this was partly due to low ZnTe:Cu transparency
and partly due to collection occurring only beyond 800 nm as expected.
ZnTe:Cu films had very low transmission above and below the bandgap
(\~560 nm) as seen in figure 18 for ZT30. Shunting and low transparency
suggested too much free Cu.

Therefore, tri-ethylamine (TEA) was investigated as a complexing agent
to reduce the chemical reactivity of the Cu in the solution. Films were
made with 21,11,6 and 1 drops of TEA. Figure 18 shows T/(1-R) measured
on glass/SnO~2~/ZnTe:Cu samples for four different TEA concentrations:
ZT51 to ZT54 have decreasing amounts of TEA but same concentration of
CuSO~4~. The films with higher TEA have much higher transmission
(transparency) at wavelengths beyond the ZnTe band gap (\~560 nm), which
signifies less defects and is an important consideration for tandem cell
development. Comparing data from ZT30 to ZT54 shows that even 1 drop of
TEA caused a significant increase in transmission below the bandgap
(\>560 nm).

![](media/image21.pct){width="3.21875in" height="3.2472222222222222in"}

[]{#__RefHeading___Toc506348168 .anchor}Figure 18 Optical transmission
measurements for ZnTe:Cu films for decreasing TEA concentration. ZT30
had no TEA. ZT51 to 54 had 21, 11, 6 and 1 drop, respectively.

Cells were made with these same 4 TEA concentrations. Cells with 21
drops behaved as though there was too little Cu, and the cell with 1
drop was unstable similar to figure 17. Cells with 11 drops were optimum
for this set of process conditions. Figure 19 shows the initial light
and dark JV curve when illuminated from front and back, respectively,
for the cell with 11 drops of TEA after heat treatment of 180°C for 10
min. There is negligible hysteresis compared to Figure 17 without TEA.
The light JV curve indicates a weakly blocking contact when illuminated
from front (through glass). The same device when illuminated from back
(through ZnTe) does not indicate a blocking contact. This indicates that
the back contact is photoconductive becoming more ohmic when
illuminated. It is not likely that this is due to differences in the CdS
photoconductivity since it is in the dark for both measurements of dark
and back illumination yet they were substantially different at forward
bias.

![](media/image22.pct){width="3.329861111111111in"
height="3.2506944444444446in"}

[]{#__RefHeading___Toc506348169 .anchor}Figure 19 JV curves for device
FS 20710F3-2 having 11 drops of TEA. Measurements for light through the
front (glass/SnO~2~) and back (ITO/ZnTe:Cu) are shown.

![](media/image23.pct){width="3.0701388888888888in"
height="2.997916666666667in"}
![](media/image24.pct){width="3.048611111111111in"
height="3.0076388888888888in"}

[]{#__RefHeading___Toc506348170 .anchor}Figure 20 Light JV curves for
back wall (left) and front wall and dark (right) illumination for FS
20710F3-4 before and after heat treatment at 180°C for 15 min.

Cells show a significant improvement after a heat treatment. Current
voltage curves measured by back wall illumination show increase in J~sc~
and V~oc~ after subsequent heat treatments. This indicates better
carrier collection. Figure 20 shows light JV curves for back wall and
front wall illumination before and after heat treatment with 21 drops of
TEA. Short circuit current went up from 1.4 mA/cm^2^ to 3.7 mA/cm^2^ for
backwall illumination. The front wall JV curves show noticeable
improvement in V~oc~ but no change in dark JV or FF or J~sc~. Heat
treatment will affect the amount of doping in the cell, as it will drive
Cu further into the film. But too much TEA will complex all the free Cu
and will not permit doping.

IEC CdS/CdTe substrates made by VT also had ZnTe:Cu contacts applied
using the same processing as above. The cells made on IEC CdTe
substrates do not show any hysteresis or blocking contact and have cell
parameters comparable to the cells from the same VT run made with
standard graphite contacts. We are working towards cells, which have
open circuit voltages over 0.80 V and higher efficiency and fill
factors.

Figures 21a and 21b show the QE with light bias measured from the front
(normal configuration) and from the back contact (allowable due to
transparent ZnTe:Cu/ITO contact). Figure 21a shows a typical QE with
little increase from 0 to --1V indicating good collection. Figure 21b
shows the response is largely limited to carriers generated very near
the CdS junction. There is no evidence of collection at a back junction,
nor is there evidence of much collection throughout the CdTe which
implies a low lifetime and low field in the bulk of the CdTe. The slight
but monotonic increase in QE with increasing depth of generation in
Figure 21a indicates increasing collection probability as carriers are
generated in the bulk of the CdTe closer to the junction. Analysis of
this region including accurate absorption coefficients may yield
information about the diffusion and drift length.

![](media/image25.pct){width="3.1104166666666666in"
height="3.001388888888889in"}
![](media/image26.pct){width="3.0506944444444444in"
height="3.001388888888889in"}

[]{#__RefHeading___Toc506348171 .anchor}Figure 21a (left) and 21b
(right). QE curves measured at --1V, 0V, and +0.5V bias with light bias
through the glass/SnO~2~ (Figure 21a) and through the ITO/ZnTe:Cu
(Figure 21b). No TEA was used.

Figure 22 shows spectral response curves for cells from a different
plate of First Solar material, which had the same chemical bath
composition but different TEA concentrations. Cells with lower TEA
showed better collection at lower wavelengths than ones with higher TEA.
Further spectral response characterization is required to fully
understand the collection mechanism for these cells.

![](media/image27.pct){width="3.5194444444444444in"
height="3.295138888888889in"}

[]{#__RefHeading___Toc506348172 .anchor}Figure 22 Backwall QE at 0V bias
for two devices with different TEA concentrations (11 and 21 drops).

Altering thickness of CdTe layers along with bifacial analysis will help
us to identify the dominant collection mechanism. We plan to do
accelerated stress testing in the near future on devices with optimized
ZnTe:Cu contacts, to provide information on contact stability. We are
working on optimizing the ZnTe deposition, heat treatments and contacts
to achieve V~oc~\>0.80 V and FF\>70% for front wall illumination and at
understanding the backwall spectral response.

## 

## **2.3 High Throughput Processing**

### 2.3.1 VT system modifications

During this year, significant design modifications were made to the
vapor transport source zone to increase CdTe utilization, improve
repeatability and improve overall robustness. The pre-deposition and
post-deposition zone heaters, which maintain zone temperature and
confine growth onto the substrate by heating the area surrounding the
source, were replaced. The original design consisted of PBN (Boralectric
Heater) disks embedded in graphite blocks. These disks are expensive and
burn out easily. The new heaters use Tantalum wire looped between a
sheet of vycor glass and a grooved boron nitride plate. These new
heaters are more robust and able to reach temperatures in excess of 700˚
C.

A new source was designed and fabricated in-house and was installed in
August 2004. The most significant change to the new source design was
the method of delivering heat to the ampoule. In the old design, a
Kanthal-A filament coiled around a quartz ampoule containing the CdTe
heated the charge directly by radiative transfer, relying on radiative
coupling between the CdTe and the filament. The carrier gas introduced
into the ampoule end transported the CdTe vapor through a series of
small holes in the ampoule into a surrounding cylindrical ceramic
manifold with a slot in the bottom directly above the substrate. The
implementation problem with this design was twofold: 1) the saturated
carrier gas passed in direct contact with the filament during transport
through the manifold and 2) the ceramic manifold's internal volume was
too great and thermal mass to great to result in insufficient heating
during a run. These problems resulted in contamination as well as
significant CdTe deposition onto the inside surfaces of the manifold.
The CdTe utilization at the substrate with this source had been 10-20%,
based on gravimetric determination of source mass loss and substrate
mass gain.[^24] The contamination resulting from filament degradation is
described below.

The new design utilizes a double-wall cylindrical boron nitride source
heater to heat the quartz ampoule containing the CdTe charge. The heater
also functions as the manifold and has a 2 mm wide slot in the bottom,
positioned over the substrate. The inner diameter of the new
heater/manifold has been reduced, from 4X the ampoule diameter to 2X the
ampoule diameter. The full specifications for the source zone are listed
in Table 9.

[]{#__RefHeading___Toc506348101 .anchor}Table 9 Specifications of IEC
vapor transport system source.

  ---------------------------- --------------------- ---------------------
  **Property**                 **Value**             **Units**

  Ampoule Material             quartz                

  Ampoule Length               15.2                  cm

  Ampoule Diameter             1                     cm

  \# Effusion Holes            12                    

  Effusion Hole Diameter                             mm

  Manifold Material            boron nitride         

  Manifold Length Overall      14.6                  cm

  Manifold Inner Diameter      1.35                  cm

  Manifold Outer Diameter      2.54                  cm

  Slit Length                  11.7                  cm

  Slit Width                   3                     mm

  Slit-to-substrate distance   1.5                   cm

  Filament Material            tantalum              

  Filament Length              3.5                   m

  Filament Diameter            0.8                   mm

  Filament Resistance          6                     Ohm

  Source Power Rating          2                     kW

  Source Current at Max P      18                    Amp

  Source T range               25-900                C

  Source gas flow rate         2-100                 sccm
  ---------------------------- --------------------- ---------------------

During operation with the old ceramic source, the manifold exterior
remained visibly dark, while the interior radiated cherry red color for
a thermocouple reading of 800ºC. With the new source, the entire
assembly is closer to an isothermal condition and is visibly uniform in
color at the same temperature reading. This evidence for nearly
isothermal condition inhibits condensation of CdTe on the inside
surfaces of the heater/manifold and results in higher utilization at the
substrate surface. Minor modifications to the source assembly also
included tightening the gas connection to the quartz ampoule and
reducing the number and diameter of exit holes in the ampoule.

The heating and cooling profiles for the source and substrate during a
run are shown in Figure 23. The deposition sequence is initiated from a
bake-out temperature of 200ºC. The substrate reaches temperature first,
after a couple of min, while resting away from the source zone. The
source reaches its set-point after 20 min, at which time the substrate
translation is initiated. After the substrate has passed through the
deposition zone, it resides under the post process heater as all
controllers are turned off. The initial cooling rate of the heater is
rapid compared to the substrate.

![](media/image28.pct){width="4.625in" height="3.9451388888888888in"}

[]{#__RefHeading___Toc506348173 .anchor}Figure 23 Heating and cooling
profiles for source (static) and substrate (moving) during a baseline
vapor transport deposition.

The average rate of CdTe mass transport out of the source is determined
by measuring the mass lost from the ampoule after each run. Likewise,
the growth rate at the substrate was determined by the mass gain. We
define a quantity, the "static equivalent growth rate", as a
figure-of-merit to compare the growth rate at different translation
rates. This is obtained by using:

$\text{StaticEquivalentGrowthRate} = \frac{\text{FilmThickness} \times \text{TranslationLength}}{\text{DepositionTime} \times \text{ZoneSize}}$,
Equation 10

where the translation length is 10 cm and the zone size is 1 cm. The
zone size is the width along the translation direction of the deposition
zone for a static deposition. The value of 1 cm was determined
experimentally to be the zone for which a uniform thickness was obtained
± 15%, for a 3 mm slit, a slit-to-substrate distance of 1.5 cm, at 20
Torr and 20 sccm flow rate onto a fixed CdS/SnO~2~/glass substrate.

In practice, deposition is initiated once the source and substrate reach
their respective set-points, by reducing the chamber pressure from \>100
Torr to operating pressure and commencing substrate translation. For a
fixed source temperature, T~sou~ \~ 800ºC, the quantity of CdTe lost
during the heating and cooling phase of the run using the differential
pressure control corresponds to 1.5 min of static deposition. Thus, for
a run with 8 min translation time, the source mass rate is computed for
8 + 1.5 = 9.5 min. For comparative purposes, the total source loss rate
and substrate gain rate are calculated in terms of moles per sec. Figure
24 shows the measured to modeled values for 8 min depositions at 20 Torr
and 20 sccm carrier gas flow rate. The source was varied from 780ºC to
840ºC, yielding a nearly 10X variation in source rate. The first-order
mass transfer reactor model used in the Figure 24 is the same used to
design the source and is described elsewhere.Error: Reference source not
found The modeled rates correspond to an assumed CdTe source slab 1 cm
wide, 10 cm long, within a cavity 1.5 cm high as the diffusion zone. The
measured data is in reasonable agreement with the model. If 100%
utilization were obtained at the substrate, the highest deposition rate
in the Figure 24 corresponds to nearly 100 μm/min static equivalent
growth rate, at the source temperature of 840ºC. At this static
equivalent growth rate, a 2 μm film could be deposited onto a 10 x 10 cm
substrate in 12 sec, while a 1 meter length (module size) could be
coated in 2 min.

Figure 25 shows the measured source rate and substrate growth rates
obtained over the same source temperature range as Figure 24 for
substrate temperature = 500ºC and at T~sou~ = 800ºC for T~sub~ = 500ºC,
550ºC and 570ºC. No growth rate dependence was found with respect to the
substrate temperature, indicating that for the range of experimental
conditions, re-sublimation of the deposited film after passing through
the deposition zone is not a critical issue. Over the range of
conditions evaluated so far, the CdTe utilization is found to be 50% ±
5%. The obtained growth rates are lower than predicted due to two
possible factors anticipated for the specific deposition conditions and
reactor: 1) absence of side shielding and 2) uncertainty in sticking
coefficients for Cd and Te~2~. In the future, these issues may be
resolved by installing shielding along each edge.

![](media/image29.pct){width="4.268055555555556in"
height="3.5416666666666665in"}

[]{#__RefHeading___Toc506348174 .anchor}Figure 24 Measured (black
squares) and modeled (red circles) source rate for 780ºC, T~sou~ \<
840ºC and 20 sccm carrier gas flow rate.

![](media/image30.pct){width="3.7944444444444443in" height="2.99375in"}

[]{#__RefHeading___Toc506348175 .anchor}Figure 25 Measured accumulation
rates at the substrate (red circles) and corresponding source rates
(black squares). Multiple samples at T~sou~ = 800ºC were deposited at
500ºC, 550ºC and 570ºC. All other samples were deposited at 500ºC.

The morphology of VTD films is somewhat correlated with growth
conditions. Over the range of conditions evaluated so far, all films
exhibit faceted morphology, with mean lateral grain size proportional to
film thickness, as shown atomic force micrographs (AFM) in Figure 26.

![](media/image31.jpeg){width="3.0395833333333333in"
height="3.1597222222222223in"}![](media/image32.jpeg){width="3.0395833333333333in"
height="3.1597222222222223in"}

[]{#__RefHeading___Toc506348176 .anchor}Figure 26 Tapping AFM images of
VT112 and VT113, deposited 2.5 cm/min and 1.25 cm/min translation speed,
yielding 3 μm and 6 μm thickness, respectively, for T~sou~ = 800ºC,
T~sub~ = 500ºC, P = 20 Torr, He carrier gas flow rate = 20 sccm, and
p(O~2~) \~ 40 mTorr.

The effect of substrate temperature for fixed 6 μm thickness is shown in
Figure 27. At temperatures, 550ºC and 570ºC, grains with lateral
dimension comparable to film thickness are found, but at 570ºC, the
average grain size is higher and the distribution of sizes is narrower.
The effect of oxygen partial pressure, which we showed earlier improves
grain packing compared to oxygen-free ambient, is not significant at
partial pressures above the baseline condition, as shown in Figure 28.

![](media/image33.png){width="3.05in"
height="3.071527777777778in"}![](media/image34.png){width="2.94375in"
height="2.7944444444444443in"}

[]{#__RefHeading___Toc506348177 .anchor}Figure 27 Tapping AFM images of
VT114 and VT120, deposited at T~sub~ = 500ºC and 570ºC, respectively,
for T~sou~ = 800ºC, 1.25 cm/min translation speed.

![](media/image35.jpeg){width="3.0006944444444446in"
height="3.1034722222222224in"}![](media/image36.jpeg){width="3.0006944444444446in"
height="3.1034722222222224in"}

[]{#__RefHeading___Toc506348178 .anchor}Figure 28 Tapping AFM images of
VT121 and VT122, deposited at p(O~2~) = 40 mTorr and 80 mTorr,
respectively, for T~sou~ = 800ºC, T~sub~ = 550ºC, and 1.25 cm/min
translation speed.

The optical characteristics of the VTD films are shown in Figures 29 and
30 for the baseline deposition condition. The optical band edge exhibits
a sharp transition at 1.50 eV. The transmission exhibits a maximum of
\~60% near the band edge and drops to 45% at 1400 nm. The high
reflectivity in the sub band gap range is a consequence of the index
mismatch between CdTe and air.

####### ![](media/image37.png){width="3.825in" height="2.4090277777777778in"}

[]{#__RefHeading___Toc506348179 .anchor}Figure 29 Optical transmission
and reflection for VTD CdTe/CdS/Ga~2~O~3~/SnO~2~/glass.

![](media/image38.png){width="3.2881944444444446in"
height="4.010416666666667in"}

[]{#__RefHeading___Toc506348180 .anchor}Figure 30 Optical absorption
versus energy for VTD CdTe/CdS/Ga~2~O~3~/SnO~2~/glass.

### 2.3.2 Impurity effects

Vapor transport deposited CdTe films are potentially sensitive to
contamination from the source heating filament and surrounding manifold.
This is because the filament and manifold must be at high temperature
and are in contact with the source vapor. Sufficient temperature is
required to radiatively couple the filament to the CdTe source material
to generate the high partial pressure needed to saturate the source
vapor. Furthermore, the source vapor will come into contact with the
manifold wall as it moves towards the emergence slit. The effects of
contamination on devices are not known but are expected to be
reactor-specific. To systematically study the effects of source
contamination, we had planned this year to employ CdTe sources of known
composition and then endeavor to correlate film composition and cell
performance with these source materials. However, over the course of
time using the ceramic source manifold, a sharp decrease in device
performance, from 11% to less than 8%, accompanied by severe shunting,
was found and appears to be correlated with quantifiable filament
degradation, which provided a different opportunity for impurity
investigation.

In a filament-heated vapor generating source, such as that used at IEC,
the selection of filament materials is based on the wire emissivity,
which determines the radiative coupling to the CdTe source material, and
the wire resistivity, which, together with the filament winding
geometry, determine the electrical power needed for heating. In the IEC
system, many metals, such as Ta, W, Ni-Cr and other alloys were
originally considered, and the alloy kanthal-A was selected, having the
composition: 70% Fe, 20% Cr, 5% Al, and 1% Si. In anticipation of
contamination by sublimation, a procedure was developed to form an oxide
diffusion barrier on the wire surface prior to installation, by heating
to \~600°C in air for 20 min.

The selection of manifold material was based on a combination of
machineable characteristics, emissivity, and cost, and Aremco
machineable ceramic was selected, composed of oxides of Si and Al.
Thermal calibration of the complete source generator showed that
operation of the kanthal filament at \~1000°C translated to a CdTe
source temperature of \~800°C and a manifold wall temperature of
\~700°C. From this, we expected and found a small CdTe deposit
accumulated on the manifold wall from run-to-run. Any contamination from
the manifold would have to diffuse through this accumulating CdTe layer.
Furthermore, measurements of the slit aperture showed no measurable
change after 80 depositions, evidence of minimal direct contamination
from the manifold.

Figure 31 shows the J-V behavior obtained before and after filament
degradation, for devices from runs VT82 and VT85. Both are the --32
piece, from the same plate location. Both devices had the same
post-deposition processing (HTA, BDH etch plus 50Å Te, then Cu/Au
contact). Figure 31 shows that cells on VT85.32 exhibited more shunting,
more hysteresis, lower V~oc~ and FF, and lower yield (1/8 vs 6/8) than
VT82.32. Figure 32 shows data from VT86.12 and a piece from First Solar
FS72435F2.11 which were co-processed at IEC, receiving the same CdCl~2~,
BDH, 50 nm Te, and Cu/Au contact. Clearly the FS device is well behaved
with no shunting or hysteresis while the VT86 device is not well
behaved. Since they had identical processing treatment, we concluded
that the origin of the poor J-V characteristics was in the CdTe layer
deposited in the IEC vapor transport system.

![](media/image39.pct){width="3.126388888888889in" height="3.15in"}

[]{#__RefHeading___Toc506348181 .anchor}Figure 31 Light and dark J-V
curves from two IEC VT samples.

![](media/image40.pct){width="3.126388888888889in" height="3.15in"}

[]{#__RefHeading___Toc506348182 .anchor}Figure 32 Light and dark J-V
curves from an IEC VT sample and a FS sample, which were co-processed at
IEC.

To determine the extent of chemical contamination, samples from
depositions made prior to and after filament degradation and after
filament replacement were analyzed by XRD, EDS and SIMS. Table 10 shows
the deposition numbers and attributes for the vapor transport system. In
the table, "new" CdTe charge indicates that the CdTe was replaced for
each deposition. "None" indicates that no CdTe powder was placed in the
source chamber. Any deposition obtained thus resulted from sublimation
from the manifold walls.

[]{#__RefHeading___Toc506348102 .anchor}Table 10 Vapor transport
deposition conditions.

  --------------------- --------------------- ----------------------------
  **Deposition**        **Filament**          **CdTe Charge**

  VT82                  Old kanthal           New

  VT85                  Old kanthal           New

  VT91                  Old kanthal           New

  VT95                  New kanthal           New

  VT96                  New kanthal           None
  --------------------- --------------------- ----------------------------

XRD and EDS analysis was conducted on film samples from these and other
runs. No differences were found in composition or diffraction patterns,
indicating no gross compositional or structural effects due to presumed
contaminants. EDS analysis was also conducted on the filament wire,
removed after deposition VT91 due to visible scaling of the filament
wire in some portions. At 30 kV, with a penetration depth exceeding 1
micron, no aluminum was detected in the wire, suggesting a loss. We also
analyzed 3 pieces of new kanthal-A wire taken at 1 meter spacings along
the wire spool and found identical composition, with 5% Al. Finally, we
had previously analyzed air-annealed wire and found similar composition,
showing that the air anneal did not appreciably change the wire surface
composition.

Sally Asher at NREL further analyzed as-deposited film samples. SIMS
depth profiles were collected with a Cameca IMS-3F using high mass
resolution for the profiles with the majority of the elements. The
profiles containing Ga and In were performed under low mass resolution
conditions since it is impossible to separate the isobaric species
^113^Cd from ^113^In, and ^115^In from ^115^Sn. All profiles were
acquired using O~2~^+^ primary ions and detection of positive secondary
ions. The data was collected from a circular area \~60 μm in diameter,
in the center of a 150x150 μm^2^ crater for ^27^Al, ^28^Si, ^32^S,
^35^Cl, ^52^Cr, ^56^Fe, ^59^Co, ^63^Cu, ^111^Cd, ^119^Sn and ^125^Te.

The Cu levels in the CdTe were found to be high for un-contacted
material, at roughly 10^18^ cm^-3^, indicating an elevated base level,
probably as a residual contaminant not removed during purification. All
of the samples exhibited Si and Al contamination. The Si contaminant
level was very high (\~0.5 at %) in sample VT91, the last deposition
with the old filament. There also appears to be significant Al
contamination in the same sample, suggesting that visible filament
scaling can transfer Al to the depositing film. The Si must originate in
the source ampoule and/or the manifold. As the source ampoules exhibited
no devitrification or erosion, we must assume that any Si originated in
the manifold.

The new boron nitride source manifold with its encased Ta filament is
chemically more inert and has already shown advantages over the original
ceramic source with Kanthal filament, which facilitates routine
fabrication of VTD cells with 11% efficiency. A systematic investigation
of CdTe source purity and contaminants on the performance of CdTe cells
is now possible and will be carried out.

### 2.3.3 High throughput post-deposition processing

Until recently, the predominant method used to treat the films involved
direct exposure of the CdTe surface to solid CdCl~2~ applied via
solution of CdCl~2~ in methanol (wet CdCl~2~)[^25] or by sublimation
from a CdCl~2~ source, either during heat treatment[^26] or prior to
heating[^27] the CdTe/CdS structure. Alternative processes by treatment
in Cl~2~ or HCl vapor have also been examined, with the result that the
CdTe surface is readily converted to CdCl~2.~[^28] The "wet" CdCl~2~
surface treatment is effective for small-cell fabrication but has
disadvantages for large area manufacture, such as non-uniform delivery
of the CdCl~2~ and handling and safe disposal of the rinsate. To
overcome these problems, vapor techniques were developed which can
quickly and uniformly deliver the chloride species to the CdTe/CdS
structure during the heat treatment without formation of post-treatment
residues.

Two atmospheric pressure approaches for delivery of CdCl~2~ vapor
species were considered: diffusion in parallel-plate configuration and
vapor transport via carrier gas in a packed bed. In both approaches, the
diffusion of CdCl~2~ through ambient gas is a critical aspect for
estimating concentration at the film surface. In a static parallel-plate
configuration, the separation is small; and diffusion calculations for
CdCl~2~ gas in N~2~ and Ar ambients show that the characteristic time to
reach equilibrium for 1 mm spacing at 400°C is a few seconds. In the
design of a packed bed delivery system, a steady state can be achieved,
but the design critically depends on the carrier gas residence time
compared to the characteristic thermal and mass diffusion time for the
source particles. Also critical to operation of any vapor CdCl~2~ system
will be the composition of the carrier gas with respect to oxygen and
water content. These components can contribute to aging of the source
thereby limiting the CdCl~2~ partial pressure and changing the chemical
reactions occurring at the CdTe surface.

### 2.3.4 Vapor chloride treatment -- diffusion delivery 

The reactor typically used for vapor CdCl~2~ treatment consists of two
parallel susceptors with independent temperature control via externally
mounted quartz halogen (FCM) lamps. This system has been used for
small-scale solar cell fabrication and to determine the solid-state
grain and grain boundary diffusion coefficients for CdS into CdTe thin
films.[^29] The lower susceptor contains solid CdCl~2~ and the upper
surface is the CdTe film undergoing the heat treatment. The primary
advantage of this reactor is the ability to decouple the CdCl~2~
concentration at the CdTe surface from the temperature of the CdTe
surface. This is achieved by maintaining the CdCl~2~ and CdTe surfaces
at different temperatures. The temperature of the CdCl~2~ surface
determines the gas phase concentration of CdCl~2~, while the temperature
of the CdTe surface determines the reaction temperature. In this way it
is possible to explore the effects of both temperature and CdCl~2~
concentration on the CdTe heat treatment process. The concentration and
temperature profiles of CdCl~2~ in the heat treatment reactor are
described as a function of *z* and *t*, where *z* is normal to the
parallel CdCl~2~ and CdTe surfaces, and *t* is time. The development of
the temperature profile is described by:[^30]

$\frac{¶\left( {\rho C}_{p}T \right)}{¶t} = \frac{¶}{¶z}\left( k\frac{¶T}{¶z} \right)$
Equation 11

where *T* = *T(z,t)* is temperature, *C~P~ = C~P~ (T)* is the heat
capacity, *k = k(T)* is the thermal conductivity, *ρ = pM/RT* is the gas
phase density, *p* is the system pressure (1 atm), *M* is the average
molecular weight of the gas phase, and *R* is the ideal gas constant.
Natural convection is minimized by maintaining the upper graphite
susceptor (touching CdTe/CdS) at the same or higher temperature than the
lower susceptor (containing CdCl~2~). There is also a time-dependent
concentration profile in the system, which is described by:

$\frac{¶c_{A}}{¶t}\text{=Ñ} \cdot \text{cD}_{\text{AB}}{Ñx}_{A}$
Equation 12

where *c~A~=c~A~(z,t)* is the concentration of component A (CdCl~2~),
and *D~AB~=D~AB~ (p,T)* is the diffusivity of component A in component B
(air), and ∇*x~A~* is the mole fraction gradient of component A. The
CdCl~2~ concentration at the source surface is specified by the
parametric vapor pressure equation and thermodynamic constants.[^31] It
is obviously a tedious process to solve the coupled Equations 11 and 12.
If it can be shown that the characteristic times for both thermal and
mass diffusion are short compared to the heat treatment time, the time
derivatives of Equations 11 and 12 may be set to zero, and a simpler
steady state solution may be obtained. The characteristic time for
thermal diffusion, *t~T~*, is given by:

$\tau_{T} = \frac{\ell^{2}{\rho C}_{p}}{k} = \frac{\ell^{2}}{\alpha}$
Equation 13

where $\ell$ is the length scale and α is referred to as the thermal
diffusivity. The characteristic time is an order of magnitude estimate
of the time required for the system to reach equilibrium. The length
scale, $\ell$, is simply the characteristic diffusion distance, in this
case the separation between the CdCl~2~ surface and the CdTe surface. We
are also concerned with the characteristic time of the concentration
profile development. The characteristic time for mass diffusion, *τm*,
is given by:

$\tau_{m} = \frac{\ell^{2}}{D_{\text{AB}}}$ Equation 14

Table 11 shows the CdCl~2~ partial pressure over the useful temperature
range, from 300°C to 500°C. Estimates of *α* and *D~AB~* for nitrogen
ambient are listed in Table 12 for the temperature range of Table 11.
The mass diffusion coefficient was determined using the Chapman-Enskog
theory, and the characteristic thermal and mass diffusive times are
shown in Figure 33 for different source to sample spacing. The time
development of CdCl~2~ concentration at the CdTe surface is shown in
Figure 34 for 400°C at the same source to sample spacing as in Figure
33. For typical source to sample spacing \~0.1 cm and treatment times of
20-30 min (1200-1800 sec), the surface CdCl~2~ vapor concentration is
found to be comparable to the equilibrium saturation pressure. Similar
results are calculated for argon and helium ambients.

[]{#__RefHeading___Toc506348103 .anchor}Table 11 Saturation pressures
(P~sat~) for materials in the CdTe-CdS-CdCl~2~ system, calculated from
published parametric thermodynamic constants \[Knacke, et al, 1997\].

+-------+-------------------+------+--------+--------+---------------+
| **T** | **P~sat~**        |      |        |        |               |
+-------+-------------------+------+--------+--------+---------------+
| **(   | **(Torr)**        |      |        |        |               |
| ºC)** |                   |      |        |        |               |
+-------+-------------------+------+--------+--------+---------------+
|       | ###### CdTe       | **   | **Te** | *      | **CdCl~2~**   |
|       |                   | Cd** |        | *CdS** |               |
+-------+-------------------+------+--------+--------+---------------+
| 300   | 2.1E-08           | 0.   | 0      | 8      | 1.3E-05       |
|       |                   | 0492 | .00038 | .4E-11 |               |
+-------+-------------------+------+--------+--------+---------------+
| 350   | 5.2E-07           | 0    | 0      | 3      | 0.00027       |
|       |                   | .294 | .00455 | .4E-09 |               |
+-------+-------------------+------+--------+--------+---------------+
| 400   | 7.9E-06           | 1    | 0      | 7      | 0.00361       |
|       |                   | .289 | .03758 | .7E-08 |               |
+-------+-------------------+------+--------+--------+---------------+
| 450   | 8.2E-05           | 4.58 | 0      | 1      | 0.0329        |
|       |                   |      | .23019 | .1E-06 |               |
+-------+-------------------+------+--------+--------+---------------+
| 500   | 6E-04             | 1    | 0.785  | 1E-05  | 0.222         |
|       |                   | 3.75 |        |        |               |
+-------+-------------------+------+--------+--------+---------------+

[]{#__RefHeading___Toc506348104 .anchor}Table 12 Thermal (*α*) and mass
(*D~AB~*) diffusivity for CdCl~2~ in nitrogen ambient.

  -------------------- -------------------- -----------------------------
  **T**                ***α***              ***D~AB~***

  **(°C)**             **(cm^2^/s)**        **(cm^2^/s)**

  300                  0.486                0.224

  350                  0.554                0.261

  400                  0.625                0.297

  450                  0.698                0.330

  500                  0.773                0.374
  -------------------- -------------------- -----------------------------

> ![](media/image41.png){width="5.028472222222222in"
> height="2.879861111111111in"}

[]{#__RefHeading___Toc506348183 .anchor}Figure 33 Characteristic time
versus temperature for thermal (dotted) and mass (solid) diffusion of
CdCl~2~ in nitrogen for different spacing.

![](media/image42.png){width="5.102777777777778in"
height="2.951388888888889in"}

[]{#__RefHeading___Toc506348184 .anchor}Figure 34 Time development of
CdCl~2~ partial pressure at CdTe surface for 400°C (D~AB~ = 0.297
cm^2^/s).

### 2.3.5 Vapor chloride treatment -- carrier gas delivery

### 

For the case of a carrier gas passing through a fixed bed of subliming
CdCl~2~ spheres, a steady state solution is needed.[^32] Mass transfer
in such a system occurs by molecular diffusion of CdCl~2~ outward from
the spherical surface into the stagnant medium and by ablation due to
momentum transfer from the carrier gas to the surface of subliming
CdCl~2~. Assuming an isothermal packed bed with preheated carrier gas,
the contribution to vapor concentration in the carrier gas along the bed
depends on bed length, and for the gas to reach saturation, the
characteristic diffusion time for CdCl~2~ must be less than the
residence time. For an isothermal bed of 0.1 cm diameter particles
operating at 400°C, with *D~AB~* = 0.297 cm^2^/s, we estimate that the
linear velocity should be less than 5 cm/s. The mass flux from each
sphere is given by:

$J_{\text{CdCl}_{2}} = k\left( C_{\text{CdCl}_{2}}^{\text{surf}} - C_{\text{CdCl}_{2}}^{\text{gas}} \right)$
Equation 15

For a subliming particle, the concentrations are determined by the
partial pressures at the surface and in the adjoining medium. The mass
transfer coefficient for a single sphere characterized by the Sherwood
and Schmidt dimensionless numbers for momentum transfer and diffusivity:

$k = \frac{D_{\text{AB}}}{d}\left\lbrack 2 + 0\text{.}6\left( \frac{\mu}{{\rho D}_{\text{AB}}} \right)^{\frac{1}{3}}\left( \frac{\text{dv}\rho}{\mu} \right)^{\frac{1}{2}} \right\rbrack$
Equation 16

where *D~AB~* is the diffusivity of the CdCl~2~ vapor in the carrier
gas, *d* is the sphere diameter, *μ* is the bulk viscosity of the
carrier gas, *ρ* is the carrier gas density, and *v* is the linear
carrier gas velocity. Due to the relatively high vapor pressure and
diffusivity of CdCl~2~, a surprisingly thin layer is required to allow
saturation pressure to be reached with \~0.1 cm particulates and 1-2
cm/s carrier gas velocity. Figure 35 shows the positional dependence of
the CdCl~2~ partial pressure within a vapor generator, which during
operation is in the vertical configuration to permit the CdCl~2~ pellets
to uniformly cover the supporting quartz frit, with the CdTe/CdS
specimen to be radiatively heated from above. The IEC apparatus utilizes
a 0.2-0.5 cm deep bed.

![](media/image43.png){width="4.699305555555555in"
height="2.620833333333333in"}

[]{#__RefHeading___Toc506348185 .anchor}Figure 35 CdCl~2~ partial
pressure versus position though a packed bed generator for operation at
400°C with 0.1 cm spheres.

### 2.3.6 Vapor chloride -- film and device results

Vapor CdCl~2~ vapor treatments yield low surface residue and uniform
device performance, making them an attractive alternative for
large-scale processing. Figure 36 shows glancing incidence x-ray
diffraction (GIXRD) patterns of the surface of as-deposited, vapor
CdCl~2~ and wet CdCl~2~ treated VT CdTe. The scans were made with Cu kα
x-ray source in asymmetric parallel-beam configuration, with incident
beam angle of 1°, which samples a depth of 120 nm into CdTe. The CdCl~2~
treatments were carried out at 420°C for 20 min. For the vapor
treatments, the CdCl~2~ sources were maintained at 400°C, which produces
a CdCl~2~ partial pressure of 3.5 mTorr. The GIXRD scan of the
as-deposited CdTe shows, in addition to the usual CdTe reflections
(111), (220), (311) and (400), a weak reflection at 2θ = 27.5° which is
a primary reflection of the native oxide TeO~2~. The vapor treated
sample exhibits the same pattern. The wet-treated sample, on the other
hand, taken after a water rinse to remove CdCl~2~, contains reflections
from non-soluble phases, indexed as CdO (2θ = 33°) and a different
crystal form of CdTeO~3~.

![](media/image44.png){width="5.117361111111111in"
height="2.776388888888889in"}

[]{#__RefHeading___Toc506348186 .anchor}Figure 36 GIXRD patterns of VT
CdTe in as-deposited and CdCl~2~ treated conditions, showing the
contrast between wet and vapor treatments.

As this work was carried out under the CdTe team activity, devices were
completed on VT material provided by First Solar. After CdCl~2~
treatment, the CdTe surface was etched and then depositing 3 nm Cu and 5
nm Au contacts. After contact processing, GIXRD scans show no detectable
oxide phases. The AM 1.5 JV performance of the best device and the
average of 8 devices made on each sample are listed in Table 13, with no
CdCl~2~ step, cells with 9% efficiency are obtainable but the
[average]{.underline} performance is much lower, due to low *V~oc~* and
FF, suggesting poor uniformity and junction quality. Higher performance
and lower standard deviation was obtained for samples processed with
CdCl~2~ treatment. Although comparable standard deviations were found
for the wet and parallel plate vapor treatments, the highest
performance, with efficiency = 12%, *V~oc~* \> 800 mV and FF \> 73%, was
obtained with parallel-plate CdCl~2~ vapor treatment. The cells made
with packed bed CdCl~2~ treatment had lower performance, with the lowest
efficiency obtained on cells at the edge, where the effluent vapor
stream mixes with ambient air. This shows that controlling the vapor
phase composition across the entire sample is important to obtaining
uniform device performance, and more optimization is necessary.

[]{#__RefHeading___Toc506348105 .anchor}Table 13 Best and average (Avg)
AM 1.5 JV results for devices fabricated from a single plate of
production vapor transport deposited CdTe/CdS and treated at 420°C using
wet CdCl~2~, parallel-plate vapor CdCl~2~, packed bed (L = 0.2 cm) vapor
CdCl~2~, and no post-deposition treatment.

+----------------+------+----+--------+----------+---------+---------+
| >              | >    |    | > **V  | > *      | >       | >       |
|  **Treatment** |  **C |    | ~oc~** | *J~sc~** |  **FF** | **Eff** |
|                | dCl~ |    | >      | >        | >       | >       |
|                | 2~** |    | > **   | > **(mA/ | >       | >       |
|                | >    |    | (mV)** | cm^2^)** | **(%)** | **(%)** |
|                | >    |    |        |          |         |         |
|                | **Pr |    |        |          |         |         |
|                | essu |    |        |          |         |         |
|                | re** |    |        |          |         |         |
|                | >    |    |        |          |         |         |
|                | >    |    |        |          |         |         |
|                |  **( |    |        |          |         |         |
|                | mTor |    |        |          |         |         |
|                | r)** |    |        |          |         |         |
+----------------+------+----+--------+----------+---------+---------+
| Vapor CdCl~2~  | 2.5  | Be | 755    | 19.1     | 64.0    | 9.2     |
|                |      | st |        |          |         |         |
| Packed Bed     |      |    |        |          |         |         |
+----------------+------+----+--------+----------+---------+---------+
|                | 2.5  | A  | 735 ±  | 19.0 ±   | 61.0 ±  | 8.5 ±   |
|                |      | vg | 20     | 0.4      | 3       | 0.6     |
+----------------+------+----+--------+----------+---------+---------+
| Vapor CdCl~2~  | 3.5  | Be | 805    | 20.3     | 73.2    | 12.0    |
|                |      | st |        |          |         |         |
| Parallel Plate |      |    |        |          |         |         |
+----------------+------+----+--------+----------+---------+---------+
|                | 3.5  | A  | 801 ±  | 20.2 ±   | 71.2 ±  | 11.6 ±  |
|                |      | vg | 5      | 0.3      | 2       | 0.4     |
+----------------+------+----+--------+----------+---------+---------+
| Wet CdCl~2~    | S    | Be | 782    | 20.1     | 68.5    | 10.8    |
|                | olid | st |        |          |         |         |
+----------------+------+----+--------+----------+---------+---------+
|                | S    | A  | 787 ±  | 20.0 ±   | 63.7 ±  | 10.0 ±  |
|                | olid | vg | 5      | 0.2      | 2       | 0.4     |
+----------------+------+----+--------+----------+---------+---------+
| None           | 0    | Be | 700    | 20.6     | 64.8    | 9.3     |
|                |      | st |        |          |         |         |
+----------------+------+----+--------+----------+---------+---------+
|                | 0    | A  | 555 ±  | 20.5 ±   | 56.5 ±  | 6.4 ±   |
|                |      | vg | 90     | 0.3      | 5       | 2.6     |
+----------------+------+----+--------+----------+---------+---------+

These results show that good device performance and uniformity can be
obtained with vapor CdCl~2~ treatment. On-going work is being focused on
reducing treatment time to \~2 min by optimizing the CdTe temperature
and CdCl~2~ partial pressure at the surface. Preliminary results with VT
CdTe deposited at IEC are shown in the next section and are very
promising.

### 2.3.7 VT cell results

Overall, the CdTe task has two principal goals: demonstrating pathways
for improved V~oc~ and for high processing throughput. The previous
sections of the report have shown progress with respect to high rate
CdTe deposition and post-deposition treatments and methods for improving
and interpreting junction operation. This section presents JV and QE
results for VT cells to evaluate the sensitivity of cell operation to
deposition and post-deposition processing, in particular the effect of
CdTe growth rate and CdCl~2~ vapor treatment. Table 14 lists the
fabrication conditions and JV parameters for selected cells in this
study. All samples were fabricated on L.O.F. (Pilkington) TEC15
soda-lime glass/SnO~2~, with 60 nm Ga~2~O~3~ HR layer and 80-90 nm CdS.
The CdTe films were deposited at 20 Torr using He carrier gas and an
O~2~ base pressure of \~ 0.02 Torr.

[]{#__RefHeading___Toc506348106 .anchor}Table 14 Best cell AM1.5 JV
results for devices fabricated with CSD CdS and VT CdTe using a single
CdTe source lot (Alfa/Aesar).

<table>
<colgroup>
<col style="width: 9%" />
<col style="width: 9%" />
<col style="width: 7%" />
<col style="width: 13%" />
<col style="width: 14%" />
<col style="width: 9%" />
<col style="width: 8%" />
<col style="width: 11%" />
<col style="width: 7%" />
<col style="width: 10%" />
</colgroup>
<tbody>
<tr class="odd">
<td><blockquote>
<p><strong>Sample</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>T<sub>dep</sub> (C)</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Growth Rate</strong></p>
<p><strong>(μm/min)</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>HTA</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>CdCl<sub>2</sub></strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Contact</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>V<sub>oc</sub> (mV)</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>J<sub>sc</sub> (mA/cm<sup>2</sup>)</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>FF (%)</strong></p>
</blockquote></td>
<td><blockquote>
<p><strong>Eff (%)</strong></p>
</blockquote></td>
</tr>
<tr class="even">
<td>118.2</td>
<td>550</td>
<td>8</td>
<td>Ar, 600C</td>
<td>410C, 20’</td>
<td>Cu/Au</td>
<td>800</td>
<td>23.6</td>
<td>70.2</td>
<td>13.2</td>
</tr>
<tr class="odd">
<td>118.4</td>
<td>550</td>
<td>8</td>
<td>None</td>
<td>410C, 20’</td>
<td>Cu/Au</td>
<td>808</td>
<td>23.8</td>
<td>69.2</td>
<td>13.3</td>
</tr>
<tr class="even">
<td>118.6</td>
<td>550</td>
<td>8</td>
<td>None</td>
<td>410C, 20’</td>
<td>Cu/C</td>
<td>745</td>
<td>23.5</td>
<td>62.1</td>
<td>10.9</td>
</tr>
<tr class="odd">
<td>119.1</td>
<td>550</td>
<td>7</td>
<td>Ar, 600C</td>
<td>410C, 20’</td>
<td>Cu/Au</td>
<td>782</td>
<td>23.4</td>
<td>69.3</td>
<td>12.6</td>
</tr>
<tr class="even">
<td>119.2</td>
<td>550</td>
<td>7</td>
<td>None</td>
<td>410C, 20’</td>
<td>Cu/Au</td>
<td>792</td>
<td>23.4</td>
<td>66.6</td>
<td>12.3</td>
</tr>
<tr class="odd">
<td>119.3</td>
<td>550</td>
<td>7</td>
<td>None</td>
<td>410C, 20’</td>
<td>Cu/Au</td>
<td>763</td>
<td>23.3</td>
<td>70.0</td>
<td>12.4</td>
</tr>
<tr class="even">
<td>119.4</td>
<td>550</td>
<td>7</td>
<td>None</td>
<td>465C, 2’</td>
<td>Cu/C</td>
<td>744</td>
<td>23.5</td>
<td>60.3</td>
<td>10.5</td>
</tr>
<tr class="odd">
<td>119.5</td>
<td>550</td>
<td>7</td>
<td>None</td>
<td>None</td>
<td>Cu/C</td>
<td>570</td>
<td>15.9</td>
<td>46.8</td>
<td>4.3</td>
</tr>
<tr class="even">
<td>120.5</td>
<td>570</td>
<td>7</td>
<td>None</td>
<td>405C, 25’</td>
<td>Cu/Au</td>
<td>794</td>
<td>23.5</td>
<td>63.4</td>
<td>11.8</td>
</tr>
<tr class="odd">
<td>120.3</td>
<td>570</td>
<td>8</td>
<td>None</td>
<td>420C, 25’</td>
<td>Cu/Au</td>
<td>763</td>
<td>24.2</td>
<td>68.6</td>
<td>12.6</td>
</tr>
<tr class="even">
<td>120.6</td>
<td>570</td>
<td>8</td>
<td>None</td>
<td>425C, 25’</td>
<td>Cu/Au</td>
<td>787</td>
<td>23.3</td>
<td>66.3</td>
<td>12.1</td>
</tr>
<tr class="odd">
<td>120.2</td>
<td>570</td>
<td>8</td>
<td>None</td>
<td>440C, 2’</td>
<td>Cu/Au</td>
<td>634</td>
<td>20.4</td>
<td>60.2</td>
<td>7.8</td>
</tr>
<tr class="even">
<td>127.4</td>
<td>550</td>
<td>12</td>
<td>None</td>
<td>410C, 30’</td>
<td>Cu/Au</td>
<td>791</td>
<td>23.6</td>
<td>68.1</td>
<td>12.6</td>
</tr>
<tr class="odd">
<td>127.1</td>
<td>550</td>
<td>12</td>
<td>None</td>
<td>415C, 30’</td>
<td>Cu/Au</td>
<td>774</td>
<td>23.2</td>
<td>63.2</td>
<td>11.4</td>
</tr>
</tbody>
</table>

The data in Table 14 allows several conclusions to be drawn about VT
cell processing. First, equivalent results were obtained with and
without a post-deposition anneal in argon, showing that the anneal step
is unnecessary, unlike PVD and electrodeposited cells, in which best
results are obtained with an annealing step prior to CdCl~2~ treatment.
Second, cells with no CdCl~2~ treatment exhibit marked photo-response,
indicating that the high deposition temperature yields an active
junction. Third, similar results were obtained over the range of static
equivalent growth rate, from 7 to 12 μm/min and deposition temperature,
from 550ºC to 570ºC. Fourth, nominally similar performance was obtained
for vapor CdCl~2~ treatment from 405ºC to 425ºC for 20-25 min but the
highest V~oc~'s were obtained over a narrower range, from 405ºC to
410ºC.

The spectral response provides information regarding final CdS
thickness, absorber band gap, and carrier collection. Figure 36 shows
the range of QE versus wavelength obtained for selected samples from
Table 14. Sample VT119.5 received NO post-deposition treatment and
exhibited the lowest overall QE, the lowest blue response and the lowest
absorber band edge = highest band gap. These observations are consistent
with no CdS diffusion into the CdTe. The other samples show a similar
range of blue response and similar long wavelength fall-off, consistent
with CdS loss ≈ 10 nm and formation of CdTe~1-x~S~x~ with x \< 0.05,
respectively. The CdS diffusion process is thermally driven and is
enhanced by grain boundary diffusion. For VT cells, having the same
large CdTe grain size, the effect is less pronounced than for PVD cells.
Sample VT118.4 was processed according to baseline procedures and
exhibited very flat response from band to band.

![](media/image45.pct){width="2.9881944444444444in"
height="3.009027777777778in"}

[]{#__RefHeading___Toc506348187 .anchor}Figure 37 QE versus wavelength
for samples from Table 14.

## **2.4 CdTe Surface Chemistry**

### 2.4.1 The mechanism of nitric acid/phosphoric acid etching of CdTe

We have previously described investigations of the surface chemistry of
CdTe during treatment with the NP etch. This etch treatment is widely
used in the community for back contact processing, forming a thick,
50-100 nm, crystalline Te surface layer. However, the chemical mechanism
of the reaction is not well understood. An understanding of the process
chemistry will allow development of more efficient and safer etch
treatments.

A 1:70:29 HNO~3~:H~3~PO~4~:H~2~O mixture was used to etch VTD CdTe
films. Films were analyzed using VASE and GIXRD. Generally, for NP
etching, as the reaction proceeds, the formation of bubbles on the CdTe
surface is observed after ∼ 20s immersion. These bubbles are due to the
formation of NO~2~(g), a product of HNO~3~ decomposition.

Due to their stability over time the NP solutions are often re-used.
When using fresh NP solutions it was observed that the etch rates of
CdTe were considerably slower than expected, with the expected bubbling
not occurring until 45 -- \>60s immersion. Subsequent etches in the same
solution showed faster bubbling times, ca.20-30s, eventually reaching a
minimum bubbling time of ∼15s. These observations suggest that a product
of the reaction is itself involved in the reaction, possibly as a
catalyst, increasing the reaction rate as the concentration of this
product increases.

To test this, and to identify the catalyst, pieces of CdTe and Te were
added to fresh NP solutions prior to etching CdTe films. This resulted
in a significant shortening of the time for bubbles to appear during
etching of CdTe films. To confirm if the catalytic species was Cd or Te
related, a small amount of Cu powder was added to a fresh NP solution.
HNO~3~ reacts with Cu similar to CdTe, with the formation of NO~2~(g).
Following addition of Cu to a fresh NP solution, the bubbling time
during etching of CdTe films decreased to ∼10s. This confirms that the
active species produced during etching is, in fact, a product of the
decomposition of HNO~3~ and not related to CdTe. Oxidation by HNO~3~ is
known to be an autocatalytic reaction, [^33] where a reaction product is
also a catalyst for the reaction. Autocatalytic processes begin very
slowly, and after an incubation period, dramatically increase in rate.
This is observed during NP etching of CdTe films by the delay in the
appearance of bubbles following, which the formation of bubbles on the
film continues rapidly.

Faster etching rates would allow shorter immersion times, which would
attenuate grain boundary etching and diffusion by the etchant, and
faster device processing times. Using CdTe, Cu or other metals to charge
etch solutions would, however, result in unnecessary and possibly
deleterious contamination. Other possible species to charge etch
solutions, preferably of similar chemical nature to HNO~3~ to avoid
possible metal contamination, were sought. The use of sodium nitrite,
NaNO~2~, which decomposes in acid to generate NO~2~(g) similar to CdTe
or Cu in HNO~3~­, to charge NP solutions was investigated. NaNO~2~ has
also been used as a catalyst for the HNO~3~ oxidation of
polysaccharides,[^34] and has been reported as a catalyst for NP etching
of CdTe,[^35] though no explanation or discussion of this was offered.
The chemical mechanics of NaNO~2~ catalysis of HNO~3~ etching are,
however, not well understood.

Care must be taken during charging of NP as the acid decomposition of
NO~2~^-^ can be violent. Bubbling of the solution following the addition
of NaNO~2~ was generally allowed to subside before beginning treatments.
For CdTe films etched in NP containing a small amount of NaNO~2~,
bubbling on the CdTe surface occurred at \<10s immersion. Bubbling times
were found to decrease further with increasing \[NO~2~^-^\], with almost
immediate bubbling observed with treatment of CdTe in NP containing
0.16M NaNO~2~. Treatment with the charged solutions results in the
formation of Te-rich CdTe surfaces as observed with the standard NP
etch.

Etch solutions of 0.05 -- 1M NaNO~2~ in 66% and 87% H~3~PO~4~:H­~2~O were
also prepared. Once the reaction between NaNO~2~ and the acid had
subsided, the resultant solutions are blue in color due to the formation
of dissolved N~2~O~3~. Treatment of CdTe, for 30s, in these solutions
showed immediate bubbling on the CdTe surface with the formation of
Te-rich surfaces. VASE measurements showed the thickest Te layers were
obtained with 0.5M NO~2~^-^ in 66% H~3~PO~4~:H~2~O and 0.05 -- 0.2M
NO~2~^-^ in 87% H~3~PO~4~:H~2~O. Higher \[NO~2~-\] resulted in
apparently thinner Te layers, due to significant etching of the
generated Te and grain boundaries by the solutions. From these results,
the use of NO~2~^-^ based etchants for CdTe back contacting appears
promising and may benefit for CdTe device processing by allowing shorter
immersion times, more efficient and controllable etching, and avoiding
the use of concentrated HNO~3~. Concentrated H~3~PO~4~, while still
required, is less toxic and corrosive than concentrated HNO~3~ and is,
therefore, easier to handle in a laboratory or industrial setting.
Optimization of the NO^2-^:H~3~PO~4~ etch and device processing are
currently underway.

Literature studies have proposed various species, including N~2~O~3~ or
NO^+^, as possible candidates for the autocatalytic species.Error:
Reference source not found^,Error:\ Reference\ source\ not\ found^ The
pre-addition of NaNO~2~, CdTe or reactive metals to the etch solution,
generates the autocatalyst through reaction with HNO~3~, which provides
fast immediate etch rates of CdTe films. The stability of the NP
solutions and the reproducibility of CdTe bubbling times indicate that
the autocatalyst is a reasonably stable species. The stability of the
blue color of the NaNO~2~/H~3~PO~4~ solutions may suggest N~2~O~3~ as
the autocatalytic species or an important reaction intermediate.
However, this has not been confirmed and N~2~O~3~ may be merely an inert
side-product formed along with the autocatalyst.

### 2.4.2 Understanding the chemistry of aniline etching of CdTe 

### 

All solutions, unless otherwise stated, were freshly prepared and
consisted of 0.2 M aniline, 0.1 M p-toluenesulfonic acid (p-TSA) and 1 M
NaCl in water. Treatments were generally carried out for 1 hr. The pH of
a standard bath is ∼6.2 and does not change following treatment. Most
samples, unless otherwise stated, consisted of ∼5 μm thick IEC VTD CdTe,
with or without CdCl~2~ treatment, generally deposited on
CdS/SnO~2~/glass substrates. Other samples tested consisted of VTD CdTe
deposited on SnO~2~/glass or just on glass. During treatments, the
samples were immersed in the solution and illuminated with a 100W bulb,
either through the substrate glass or directly on the CdTe. Solution
temperature was not controlled during treatment, however, this did not
rise above mid-30°C's with illumination. Dark treatments were carried
out by completely wrapping the beaker in aluminum foil and storing in a
dark cupboard. The surfaces of the treated CdTe films were characterized
using VASE and GIXRD.

As observed by the Toledo group, following a successful treatment, the
CdTe surface changed color to light gray, though it is not always
uniform (see later). VASE and GIXRD both confirm the presence of
crystalline elemental tellurium on the CdTe surface following successful
treatments under illumination. Figure 38 shows the VASE ε~2~ vs. photon
energy spectra of a CdTe film before and after illuminated treatment in
a standard bath. Before treatment the spectrum contains two peaks due to
CdTe, however, after etching the spectrum now contains a single peak ∼2
eV due to Te.[^36]^,^[^37] The thickness of the generated Te film has
not been modeled, however, the CdTe peaks cannot be discerned in the
VASE spectrum of the treated sample, suggesting that a Te film of 50-100
nm is formed. Figure 39 shows the GIXRD pattern, of a CdTe/CdS sample
following illuminated treatment in a standard bath. The Te(101) peak is
present at 27.6°, while the peak at 26.5° can be assigned to CdS exposed
at the edge of the piece. Similar results were also obtained when the
CdTe film was directly illuminated. Longer treatments resulted in the
formation of a greater amount of surface Te. Figure 40 shows the GIXRD
pattern of a CdTe/CdS sample that had received a 5.5 hrs illumination
through the substrate glass in the standard bath and shows more Te has
been generated on the CdTe surface compared 1 hr treatments. A dark
organic film, most likely polyaniline, was visible on the surface of
this sample. VASE spectra of this piece were very similar in intensity
to the spectrum shown in Figure 38, despite a greater amount of Te being
generated. A stronger Te peak is expected, however, the presence of the
organic film may result in a decrease in the VASE spectral intensity.
AFM measurements have also detected the presence of an organic or
polymer-like species on the surface of treated samples. The presence of
a Te-rich back surface will result in an improvement of devices
contacted with a Cu-based back contacts and account for the performance
improvement observed by UT between treated and non-treated cells. The
effect, if any, of adsorbed polyaniline on device performance is not
clear.

![](media/image46.png){width="3.459722222222222in"
height="2.9916666666666667in"}

[]{#__RefHeading___Toc506348188 .anchor}Figure 38 VASE spectra of CdTe
film (a) before and (b) following 1 hr treatment in a standard aniline
bath with illumination through the substrate glass.

![](media/image47.png){width="3.7805555555555554in"
height="2.998611111111111in"}

[]{#__RefHeading___Toc506348189 .anchor}Figure 39 GIXRD, at 1°
incidence, of CdTe film following 1 hr treatment in a standard aniline
bath with illumination through the substrate glass.

![](media/image48.png){width="3.78125in" height="2.998611111111111in"}

[]{#__RefHeading___Toc506348190 .anchor}Figure 40 GIXRD, at 1°
incidence, of CdTe film following 5.5h treatment in a standard aniline
bath with illumination through the substrate glass.

Some treatments were carried out under just ambient laboratory lighting
at room temperature. After a 1 hr treatment, no visible change could be
seen on the CdTe surface. VASE, however, showed some subtle changes,
consisting of a slight decrease in the intensity of CdTe peaks and the
formation of a weak peak around 2.5-3 eV. The new peak, due to its high
energy, is unlikely to originate from Te. Preliminary modeling of the
data suggests that these changes may be due to the presence of organics
on the CdTe surface, possibly adsorbed aniline. This indicates that the
system is very sensitive to the light, with only ambient lighting
sufficient to begin the reaction by allowing surface adsorption of
aniline. VASE results from an illuminated treatment of CdTe, where the
sample was removed at various periods and measured before continuing
treatment, indicate that the initial stages appear to involve an
adsorption process, which attenuates the CdTe spectrum. After a period
of time, generation of Te begins. These initial changes were not
observed for treatments in aniline-free baths, indicating that aniline
is the surface-active species. This adsorption step is also not observed
with treatments in complete darkness, suggesting that aniline adsorption
may be due to surface charge generated with illumination. For some
treatments, the aniline was substituted with NH~3~ or NH~4~Cl, which
have similar acid-base behavior, but different redox chemistry, to
aniline. No reaction occurred with the CdTe surface in either of these
baths.

Illuminated treatments were carried out using baths containing varying
aniline concentrations. Baths of high concentrations, 0.4 M, were
cloudy, due to the low solubility of aniline at this pH. However, the
VASE of the treated CdTe surface was very similar to those obtained
following treatment in standard baths. Treatment in baths of lower
aniline concentrations, 0.1 M, showed practically no reaction over the
standard 1 hr treatment time. Experiments carried out in the absence of
aniline, with the bath containing only p-TSA and NaCl, showed no surface
color change and only subtle changes to the surface were observed with
VASE, similar to those with ambient room light during treatment. These
observations confirm the direct involvement of aniline in the reaction
to form Te.

Increasing the p-TSA concentration to 0.06 M (pH = 5.8) resulted in no
surface reaction on the CdTe, which is consistent with the chemistry of
aniline, [^38] where the protonated form of aniline is harder to
oxidize. This also confirms that the formation of Te is directly related
to the reactivity of aniline. Also, the acid may be competing with, or
blocking, the adsorption of aniline to the CdTe leading to an
attenuation of the bath reaction. When acid is not added to the bath,
i.e. treatment is carried out in aqueous aniline-NaCl solution, the
surface reaction does not occur. A set of experiments were carried out
where the p-TSA was substituted with a range of different acids,
including oxalic, maleic, phthalic, ascorbic, acetic, phosphoric and
hydrochloric acids and EDTA. The pH of each bath was from 5.0 -- 6.0,
each bath produced a Te-rich CdTe surface, similar to the standard bath.
However, for the HCl bath, some etching of the surface may have been due
to the HCl itself. These results indicate that the nature of the acid is
not critical to bath behavior, though acid concentration and control of
bath pH is important.

Varying NaCl concentration was also found to have a significant effect
on the surface reaction. Carrying out illuminated treatments in baths
containing decreased NaCl concentration, 0.5 M, significantly slowed the
reaction rate and only a very small amount of Te was formed on the
sample surface. In baths where NaCl was not added, no surface reaction
occurred. Doubling the NaCl concentration resulted in a significant
surface reaction as well as the formation of a visible organic film on
the CdTe surface. It has been previously shown that the choice of
electrolyte can have an effect of the oxidation chemistry of
aniline.Error: Reference source not found The NaCl was substituted with
other electrolytes, including KCl, NaBr and NaI. For an illuminated
treatment, with the bath containing 1 M NaBr, no surface reaction
occurred, but with 1M KCl a similar surface reaction to that of the
standard baths was observed. However, using 1M NaI as the electrolyte
led to a slow surface reaction with CdTe, resulting in a significantly
thinner Te film. From these results, it appears that the presence of
Cl^-^ ions is critical for the reaction to occur at a significant rate.

As the bath temperature does rise slightly during illuminated
treatments, the role of temperature in the reaction was investigated. A
CdTe sample was treated in the dark in a standard bath held at 40-50°C
for 1 hr. From VASE measurements, no changes to the surface occurred.
Following treatment in a standard bath at ∼90°C in the dark, a
significant surface reaction occurred. The surface became uniformly
light gray, and a thin dark organic film was visible on the surface.
VASE indicated the formation of significant surface Te, however, the
spectral intensity was very low due to the presence of the organic film.
Later experiments found that the formation of Te at high temperature and
in the dark occurred independent of the presence of aniline, requiring
only the presence of aqueous NaCl and p-TSA. This is an interesting etch
treatment for CdTe, which requires further investigation for it
applicability to device processing.

The non-uniformity sometimes observed for the surface reaction during
illuminated treatments may be due to un-dissolved aniline droplets
attaching to the surface on immersion of the sample, which can block the
surface reaction. Sometimes the surface reaction does not occur near the
clip holding the sample and may be due to effects of shadowing. This was
confirmed by carrying out a treatment with a CdTe substrate, which had a
small cross pattern of tape placed on the glass. After a 90 min
treatment in a standard bath with illumination through the substrate
glass, a pattern of unreacted CdTe in the shape of the tape cross
pattern was observed on the film surface. VASE confirmed that surface Te
was generated only outside the taped area. This confirms the
photo-catalytic nature of the reaction and that only CdTe over areas of
the sample that have received direct illumination are involved in the
reaction.

Carrying out treatments with non-CdCl~2~ treated CdTe resulted in the
formation of Te-rich surface layers, though the rate of the reaction was
slower than observed for CdCl­~2~-treated films. For example, a 90 min
aniline etch of non-CdCl~2~ treated pieces produced similar Te-levels to
CdCl~2~-treated films following 60 min bath treatments. This may be due
to morphological effects or lower conductivity of non-treated films. To
investigate the role, if any, of the other substrate layers during
aniline treatment of CdTe, different substrates; CdTe/SnO~2~/glass (no
CdS) and CdTe/glass (no CdS or TCO), were tested. Treatments were
carried out with illumination on either the glass or the CdTe side for
90 min. In each case, a Te-rich CdTe surface was produced. These results
indicate that only the CdTe layer is involved in the reaction during
aniline treatment under illumination. Treatment of glass slides, either
in light or in the dark at 90°C, resulted in the visible adsorption of
organic species, most likely polyaniline, on the glass surface. This
indicates that the polymerization of aniline can be caused by photo and
thermal treatments of the bath solution and that the polymerization is
not necessarily involved in the CdTe→Te surface reaction.

The involvement of O~2~ in the surface reaction has also been confirmed.
Treatments of CdTe carried out in Ar(g) purged baths show a significant
decrease in the rate of Te generation. However, some Te was still
produced, likely due to the presence of O~2~ adsorbed on the CdTe
surface. This was tested by removing adsorbed O~2~ from the CdTe
surface, by 20-30 min treatments at 190°C in an Ar(g) atmosphere, before
transporting the piece under Ar(g) and carrying out treatment in an
Ar(g)-purged aniline solution. An Ar(g) blanket was maintained over the
solution throughout the 90 min treatment. No Te was formed on the CdTe
surface, confirmed by VASE and GIXRD. Storing the same solution
overnight and then repeating the treatment with a standard CdTe sample
further tested this observation. This time the treatment resulted in the
formation of a Te-rich surface. The formation of Te on the second
treatment does confirm the critical role of O~2~ in the treatment of
CdTe with aniline solutions. The confirmation of the involvement of O~2~
in the surface reaction accounts for a reduction half-cell for the
surface reaction. O~2~ can be reduced by conduction band electrons and
will complement the oxidation of aniline and/or CdTe.

We have found that, unlike for other etches, the generated Te is
unreactive towards the etch solution. For example, treating evaporated
Te films or Te-rich CdTe surfaces in the aniline solutions, either under
illumination or at 90°C in the dark, showed almost no reaction. The
stability of Te was further confirmed by monitoring the time dependence
of a treatment of CdTe in a standard aniline solution. The sample was
removed at various times from the bath and analyzed using VASE, before
returning to the bath and continuing treatment. With time, Te begins to
be formed, slowly reaching a maximum after ∼400 min immersion. However,
with further treatment there was little change in the VASE spectra,
highlighting the stability of the generated Te and the self-limiting
nature of the reaction. When the surface becomes saturated with Te, the
reaction ceases and, thus, protecting the newly formed Te-rich CdTe
surface.

The University of Toledo has reported that devices treated with the
aniline bath exhibit very good uniformity of initial performance,
compared to those treated with Br~2~- or HNO~3~-based etches. We
attribute this to the reproducibility and stability of the Te-rich
surface during treatment. While this treatment may be slower compared to
other etches, the self-limiting nature of the reaction and the stability
of the generated Te prevent the CdTe surface from being over-treated, as
can occur during Br~2~ and NP etching.

We can conclude that a successful aniline treatment of CdTe results in
the formation of a crystalline Te-rich surface layer on the CdTe
surface, which will be beneficial for back contact processing. A
successful treatment requires the presence of aniline, Cl^-^ ions, O~2~,
consistent illumination and careful control of pH. Long or fast
treatments also result in the deposition of an organic film, likely
polyaniline, on the surface, though this step has been shown to occur
with just illumination or heating and has not been confirmed to be
involved in the CdTe surface reaction.

Figure 41 shows a schematic of the proposed chemical mechanism for the
aniline treatment. The first consideration is the nature of the CdTe
film. The reaction occurs irrespective of the substrate and
CdCl~2~-treatment. This does suggest that the CdTe film does not behave
as a solar cell or as a single crystal semiconductor electrode, but more
like a colloidal semiconductor solution at these conditions and can be
imagined as an immobilized colloidal suspension. This has implications
on the reactivity of the film, in that both holes and conduction band
electrons generated by illumination can both diffuse to the grain
surface and become involved in the reaction. Similar mechanisms are
proposed for photocatalysis of organics in wastewater using TiO~2~
colloidal suspensions.[^39]

Aniline is possibly adsorbed on the film surface, which will allow
easier reaction with the CdTe. This adsorption may be photo-assisted.
Aniline can be oxidized by the generated holes to either form
polyaniline, which may adsorb to the CdTe surface as well, or complete
mineralization to CO~2~. From literature,Error: Reference source not
found the presence of certain anions affects the oxidation of aniline
and similar molecules. The presence of Cl^-^ salts has been shown to
decrease the aniline oxidation potential, Error: Reference source not
found therefore, making it easier to oxidize and would allow the etch
reaction to occur at a faster rate. A possible reason for this may be
due to complexation between the Cl^-^ and aniline, or its protonated
form, destabilizes the aniline structure and becomes easier to oxidize.

Complementing the oxidation of aniline is the reduction of O~2~.
Hydrogen peroxide, H~2~O~2~, superoxide ions, O~2~^-^, or hydroxide
radicals, OH**^.^**, have been proposed as products of O~2~ reduction
under such conditions.Error: Reference source not found These species
are all strong oxidizing agents and have been proposed as the active
agents for the removal of organics during photocatalysis of wastewater
with colloidal TiO~2~. We propose that these generated species oxidize
the CdTe surface to form crystalline Te. When the surface becomes
saturated with newly generated Te, the self-limiting photocatalytic
reaction ceases. This both protects the Te layer, as no further
oxidizing agents can be generated and gives very reproducible surfaces.

[]{#_Ref83193172
.anchor}![](media/image49.png){width="3.745138888888889in"
height="3.1444444444444444in"}

Figure 41 Schematic of the proposed chemical mechanism for the formation
of surface Te during aniline-based etching of CdTe under illumination.

Devices have been processed by treating CdTe/CdS/TCO/glass substrates
for 60 min in standard aniline solutions under illumination and
completed with Cu-containing back contacts and graphite paste. Initial
devices show very promising performance, with V~oc~ of ∼800 mV, J~sc~
∼20 mA cm^-2^, and fill factor ∼60% or higher. Illuminating the
substrates either through the glass or directly on the CdTe was found to
have little effect on device performance. The effect of different
treatment times and device stability will be further investigated.

# 3 CuInSe~2~-based Solar Cells

## **3.1 In-line Deposition System Improvements:**

## 

### 3.1.1 Cu source

Two aspects of the Cu sources have been addressed. The first one
concerned the spitting of Cu droplets from the nozzles resulting in
inclusions in the CIGS films causing shorting of the devices. The issue
was resolved by changing the shape of the nozzles from cylindrical to
conical reducing the temperature drop along the nozzle. The details are
described in appendix 1, which is a copy of the paper presented at the
Solar Energy Progress Review Meeting in October 2004. Though Ga and In
sources have not showed any spitting, the conical nozzle design was
incorporated into these sources as well.

A major reliability issue in the in-line system has been the operational
life of the Cu source. After about ten to twenty runs, cracks developed
on the BN internal boat at the thermocouple well causing liquid Cu to
leak onto the thermocouple. This gradually destroys the thermocouple
(TC) but because of the Cu solidifying in the TC well the boat and the
TC had to be replaced. In order to avoid this problem, a new Cu source
was designed that avoided sharp internal lines and corners as they
invariably act as stress concentrators. The resulting boat, which also
was designed to have a larger capacity, has been put into service and
has been operating reliably for some time now. Since, the issue of
liquid Cu leaking into the TC well has been resolved, a naked C type
thermocouple, rather than a sheathed one is now being used.

In addition to the design improvements a new method was developed to
improve the accuracy of determining actual melt temperature in the Cu
source. As the thermocouple ages a non-negligible difference between TC
read-out and actual melt temperature was observed by monitoring the
solidification point during the cool-down. In order to quantify this
\"offset\" and correct the set-point, before each run the Cu boat is
heated past the melting point and then cooled. Temperature as a function
of time, T(t), is recorded during cool down. In the first derivative,
dT/dt, the solidification phase transition shows up as a distinct peak,
and in the second derivative, d^2^T/dt^2^, as a zero crossing. Figure 42
gives an extreme example where, taking the melting point of Cu as
1085°C, an offset of --167°C is observed.

![](media/image50.jpeg){width="5.026388888888889in" height="3.0in"}

[]{#__RefHeading___Toc506348192 .anchor}Figure 42 Determining the TC
offset for the Cu source during a run.

### 3.1.2 Se delivery

The new Se manifold was evaluated with a test matrix of 2 and 4 effusion
holes having diameters of 3/32\" and 3/64\". The tests showed that a
3/8\" manifold with two 3/64\" holes placed at either side of the Ga
source with the source at 300°C was acceptable. Under this particular
condition 14 runs were made. The compositional analysis of the CIGS
films as obtained by EDS is given in Table 15.

The data in Table 15 show acceptable uniformity both run-to-run and
within a longer run, i.e. run No.: 70265 where sampling points were 9\"
apart.

However, by the end of this series of experiments it was found that
graphite float keeping the thermocouple on the surface of the liquid Se
was soaked with Se due to its porosity. The result of which was that
temperature did not represent true surface temperature. A larger Se
source was put in service where along with a new graphite float for the
surface temperature there is also a fixed thermocouple to measure the Se
temperature at the bottom of the source. The goal is to be able to
control precisely the Se delivery.

[]{#__RefHeading___Toc506348107 .anchor}Table 15 Composition, by EDS, of
CIGS films over 14 runs.

  ------------ ----------- ----------- ----------- ----------- ------------- -------------
  **Run \#**   **Cu**      **In**      **Ga**      **Se**      **Cu /**      **Ga /**

  ** **        **(at%)**   **(at%)**   **(at%)**   **(at%)**   **(Ga+In)**   **(Ga+In)**

  70261.12     24.2        17.9        8.7         49.3        0.91          0.33

  70262.22     22.0        18.2        8.2         51.6        0.83          0.31

  70264.22     24.0        18.5        8.0         49.6        0.91          0.30

                                                                             

  70265.12     24.9        18.6        8.4         48.1        0.92          0.31

  70265.22     24.4        18.7        7.6         49.3        0.93          0.29

  70265.32     23.3        18.8        8.0         50.0        0.87          0.30

  70265.42     23.6        18.6        8.4         49.4        0.88          0.31

  70265.52     23.2        18.6        7.9         50.2        0.87          0.30

  70265.62     22.9        18.2        9.1         49.9        0.84          0.33

  70265.72     24.4        15.9        10.4        49.4        0.93          0.39

  *Average*    *23.8*      *18.2*      *8.5*       *49.5*      *0.89*        *0.32*

  *Std Dev*    *0.8*       *1.0*       *0.9*       *0.7*       *0.04*        *0.04*

                                                                             

  70268.12     22.3        20.0        7.3         50.4        0.82          0.27

  70268.32     22.1        19.6        7.9         50.4        0.80          0.29

  70269.22     21.5        19.5        6.9         52.1        0.81          0.26

  70270.22     24.0        17.9        7.9         50.2        0.93          0.31

  70271.22     24.0        18.7        7.7         49.7        0.91          0.29

  70272.12     21.6        19.9        7.6         50.9        0.79          0.28

  70272.22     23.2        19.4        7.7         49.7        0.86          0.28

  70275.22     23.6        17.8        8.1         50.6        0.91          0.31

  70276.22     24.0        17.8        8.1         50.2        0.93          0.31

  70277.22     23.6        18.3        7.6         50.5        0.91          0.29

  70278.22     24.7        18.1        7.7         49.5        0.96          0.30

  70279.22     24.4        18.3        8.3         49.1        0.92          0.31

  *Average*    *23.3*      *18.6*      *7.9*       *50.2*      *0.88*        *0.30*

  *Std Dev*    *1.1*       *0.8*       *0.4*       *0.8*       *0.05*        *0.02*
  ------------ ----------- ----------- ----------- ----------- ------------- -------------

### 3.1.3 Web temperature control

A new web shield with thermocouple holders attached to it has been put
in place. These thermocouples are sheathed ungrounded K-type of 0.010\"
outside diameters and are positioned to touch the web\'s coating side by
their natural \"spring\" action just before and just after the
deposition zone. For heater temperatures of 500°C and 550°C over
respectively the first half and the second half of the deposition zone,
the respective web temperatures observed were around 400°C and 420°C.
These thermocouples are now used as control thermocouples for the two
substrate heaters.

### 3.1.4 Film cracking

In the course of developing more reproducible precise control of Se
delivery it was observed that the extent of film cracking depended on
the Se flux. More precisely, a dramatic decrease in the crack density
was observed when Se flux was reduced by a factor of 10, while still
sufficient to form CIGS. It was postulated that the mechanism of
cracking was the chemical reaction between Mo and Se. This reaction
would reduce the yield strength of Mo causing it to crack under the
thermal and mechanical stresses the web is subjected to in the
roll-to-roll coating process. In order to test this hypothesis, Mo films
were deposited with oxygen introduced into the film. The idea was based
on the fact that molybdenum oxides are chemically more stable than
selenides. Appendix 1 gives the detail of the results obtained with
oxygenated vs non-oxygenated Mo films. Mo films containing 8 at% oxygen
showed no cracking while Mo films with around 3 at% oxygen displayed
heavy cracking. It can be concluded that the minimum oxygen level in the
Mo film is somewhere between 3 at% to 8 at%. It should be noted that
oxygen need not be distributed uniformly throughout the Mo film. Since
the idea is to eliminate the reaction with Se, only the top portion of
the Mo film needs to contain oxygen. In fact Mo film prepared by adding
oxygen into the sputtering gas discussed in Appendix 1 had around 8 at%
oxygen in the top 1000Å of the film while the first 1000Å contained only
3 at%.

## **3.2 CuInSe~2~-based devices with increased 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~). In addition to research on Cu(InGa)Se~2~ and CuIn(SeS)~2~ alloys
at several laboratories worldwide, IEC has previously reported on thin
film solar cells using Cu(InAl)Se~2~ [^40]^,^[^41]^,^[^42] 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 smaller change in lattice parameter.

Cu(InAl)Se~2~ films have been 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.Error: Reference source not found 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.Error: Reference source not found However,
efforts to study devices with increasing bandgap or to implement
different deposition processes were still severely hampered by poor
Mo/Cu(InAl)Se~2~ adhesion resulting in many films delaminating during
device processing or characterization. An inverse relation between the
Al content at the back of the Cu(InAl)Se~2~ film and adhesion was
observed. Consequently, efforts on Cu(InAl)Se~2~ were focused during
this period on obtaining a more fundamental understanding of the
Mo/CuInSe~2~-alloy interface.

In addition, we have previously identified poor adhesion at the Mo
interface as a critical issue in the fabrication of cells using inline
evaporation of Cu(InGa)Se~2~ on flexible polyimide substrates [^43] and
in this case the adhesion was again improved by the deposition of excess
Ga at the Mo interface.

The present effort is focused on formation of a MoSe~2~ interfacial
layer during the CuInSe~2~-alloy deposition.[^44] The role of MoSe~2~ in
Cu(InGa)Se~2~ devices is not well understood, even though it plays a
role in the formation of ohmic contact between Mo and Cu(InGa)Se~2~.
However, MoSe~2~ has a layered structure, similar to mica, in which
planes with Se-Se bonds can be easily cleaved leading to adhesion
problems. TEM studies by Nishiwaki et. al.Error: Reference source not
found showed that the orientation of the MoSe~2~ plays an important role
in controlling the adhesion properties. Adhesion is best with the
MoSe~2~ c-axis parallel to the Mo layer so the Se-Se planes are
perpendicular to the substrate. In this case, an XRD pattern will show
(100) and (110) diffraction peaks. Delamination is more likely with the
Se-Se planes parallel to the substrate; in which case, (002) diffraction
peaks will be observed. Thus, a specific objective of this work is to
determine how Ga and/or Al at the Mo interface affects the MoSe~2~
formation and orientation. Additional objectives include determining if
growth conditions for the Mo deposition or Cu(InGa)Se~2~ evaporation can
control the MoSe~2~ formation, and how it is affected by the presence of
Na from the glass substrate.

Experiments were done to characterize the formation of MoSe~2~ by
reaction of Mo layers, deposited on either soda lime (SLG) or
borosilicate glass (BSG), which is Na free, with evaporated Se. In
addition, 20nm thick Ga or Cu layers were deposited on Mo films prior to
Se reaction, which was done at 500°C for 1 hr. Al was not included in
these initial experiments because of the stability of Al~2~O~3,~ which
would form. Films were characterized by XRD to determine the phases
formed by the selenium reaction and their orientation. XRD measurement
using Bragg-Brentano geometry was carried out from 10° to 80°. This
configuration enforces θ-2θ symmetry so only lattice planes parallel to
the substrate planes would satisfy the Bragg equation resulting in
diffraction peaks. Glancing incidence XRD (GIXRD) measurements were done
at 0.3°, 0.4°, 0.5°, 0.6°, 0.8° and 1.0° incident angles, all above the
critical angle, which correspond to a sampling depth varying from 80 nm
to 270 nm. In the GIXRD measurements, the incident angle is fixed and
the reflections satisfying the Bragg conditions are obtained from
lattice planes tilted at different angles to the sample plane. If
measurements are carried out at a different incident angle, a different
set of lattice planes with the same spacing but tilted from the sample
plane satisfy the Bragg equation. In the present set of experiments, the
variation in incident angle is sufficiently small, that the same set of
lattice-planes, satisfy the Bragg reflection conditions.

XRD and GIXRD scans are shown in Figures 43 for selenized BSG/Mo,
SLG/Mo, BSG/Mo/Ga, SLG/Mo/Ga, BSG/Mo/Cu, and SLG/Mo/Cu samples. The
GIXRD data is shown with increasing incident angle displaced to higher
intensity. For all samples, the XRD spectra show MoSe~2~ peaks, which
are labeled by their (hkl) indices. Mo, Ga~2~Se~3~ and Cu~2~Se~3~ peaks,
when observed, are labeled by their phase but not indices. The (103)
peak for hexagonal MoSe~2~ has the highest structure factor and
therefore intensity in powder patterns but is absent in all the films.
This is consistent with previous results reported for MoSe~2~ films
formed by selenium reaction of Mo films [^45]. Only (002) reflections,
corresponding to the MoSe~2~ layered structure parallel to the substrate
plane, and (100) and (110) reflections, arising from layers
perpendicular to the substrate plane, are observed.

The orientation parameter for the three MoSe~2~ peaks was determined
from the XRD peak heights using the Harris method in which a value \> 1
for a given peak indicates preferred orientation compared to a random
powder diffraction pattern. The results are listed in Table 16. This
shows that the MoSe~2~ films have a preferred orientation in the (110)
direction, for all cases. The films with the Cu layer, on either type of
glass, have a stronger relative (002) orientation than the selenized Mo
or Mo/Ga layers. There were no significant differences between the BSG
and SLG substrates.

The GIXRD spectra, shown at multiple angles, have much less noise since
they sample a much greater volume of the phases near the surface. In
general, all the diffraction peaks in the symmetric XRD spectra are also
observed in the asymmetric GIXRD spectra but with different relative
peak heights. For the MoSe~2~ phase, the most significant difference is
the (110)-peak height, which is the largest peak in the XRD spectra; but
is much smaller in the GIXRD data. The symmetric XRD measurements
indicate that (110) lattice planes lying parallel to the sample are
preferred. However, the asymmetric GIXRD measurement requires (110)
planes to be tilted by \~27^o^ from the sample planes to satisfy the
Bragg reflection condition so the peak is much smaller.

[]{#__RefHeading___Toc506348108 .anchor}Table 16 Orientation parameter
for different peaks determined using the Harris method with 3 (hkl)
peaks.

  --------------- ----------------- ------------ ------------ ---------------
                  **Structure**     **P(002)**   **P(100)**   **P(110)**

  IE0701          7059/Mo           0.3          1.1          1.7

  IE0702          7059/Mo/Ga        0.2          0.9          1.9

  IE0703          7059/Mo/Cu        0.6          0.9          1.5

  IE0704          SLG/Mo            0.2          1.1          1.7

  IE0706          SLG/Mo/Ga         0.2          0.8          2.0

  IE0708          SLG/Mo/Cu         0.6          0.7          1.6
  --------------- ----------------- ------------ ------------ ---------------

Additional observations from the GIXRD data include:

-   For the selenized Mo, the ratio of (100) / (002) peak heights
    increased with increasing incident angle.

-   For selenized Mo/Ga films the relative (002) and (110) peak heights
    behave similarly to the Mo films but the (002) peak is significantly
    smaller, especially on the SLG substrate. Ga~2~Se~3~ peaks are also
    observed.

-   For the Mo/Cu on SLG and BSG substrates, the largest MoSe~2~ peak is
    the (002). The relative (002) and (100) peak heights do not change
    with incident angle. Instead, Cu~2~Se~3~ and both (002) and (100)
    peaks increased with incident angle.

(a)![](media/image51.png){width="3.928472222222222in"
height="2.9583333333333335in"}

(b)![](media/image52.png){width="3.972916666666667in"
height="2.967361111111111in"}

[]{#__RefHeading___Toc506348193 .anchor}Figure 43 XRD (a) and GIXRD (b)
spectra from selenized BSG /Mo substrates.

(a)![](media/image53.png){width="4.11875in"
height="3.097916666666667in"}

(b)![](media/image54.png){width="4.11875in"
height="3.097916666666667in"}

[]{#__RefHeading___Toc506348194 .anchor}Figure 44 XRD (a) and GIXRD (b)
spectra from selenized SLG/Mo substrates.

(a)![](media/image55.png){width="4.190972222222222in"
height="3.1493055555555554in"}

(b)![](media/image56.png){width="4.190972222222222in"
height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348195 .anchor}Figure 45 XRD (a) and GIXRD (b)
spectra from selenized BSG /Mo/Ga substrates.

(a)![](media/image57.png){width="4.190972222222222in"
height="3.1493055555555554in"}

(b)![](media/image58.png){width="4.190972222222222in"
height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348196 .anchor}Figure 46 XRD (a) and GIXRD (b)
spectra from selenized SLG/Mo/Ga substrates.

(a)![](media/image59.png){width="4.190972222222222in"
height="3.1493055555555554in"}

(b)![](media/image60.png){width="4.190972222222222in"
height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348197 .anchor}Figure 47 XRD (a) and GIXRD (b)
spectra from selenized BSG/Mo/Cu substrates.

(a)![](media/image61.png){width="4.190972222222222in"
height="3.1493055555555554in"}

(b)![](media/image62.png){width="4.190972222222222in"
height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348198 .anchor}Figure 48 XRD (a) and GIXRD (b)
spectra from selenized SLG/Mo/Cu substrates.

To characterize the MoSe~2~ formation in the presence of Al as in the
case of Cu(InAl)Se~2~ deposition, Mo/Al/Cu layers were deposited to
minimize oxidation of the Al. After Se reaction, Cu-Al intermetallic
alloys were identified by XRD measurements but formation of MoSe~2~ was
negligible. Thus, this structure may not be suitable for
characterization of the MoSe~2~ growth and orientation with Al alloys.
An alternative method is to peel the CuInSe~2~-alloy film off of the
substrate using epoxy. With Cu(InGa)Se~2~ films, this was found to
cleanly separate the Cu(InGa)Se~2~ from the Mo with no Mo detected on
the backside of the Cu(InGa)Se~2~ film.

Figure 49 shows a symmetric θ-2θ XRD scan of the Mo/MoSe~2~ remaining on
the substrate after peeling a Cu(InGa)Se~2~ film. The Cu(InGa)Se~2~ in
this case was deposited by elemental evaporation at 550°C. The peaks
identified by number in Figure 49 are listed in Table 17 and indexed as
either Mo, MoSe~2~, or Cu(InGa)Se~2~. Thus, there is some residual
Cu(InGa)Se~2~ on the substrate. The MoSe~2~ is oriented in the (002)
direction indicating Se-Se layers parallel to the substrate.

![](media/image63.pct){width="5.059722222222222in"
height="2.747916666666667in"}

[]{#__RefHeading___Toc506348199 .anchor}Figure 49 XRD scan of the
glass/Mo substrate after a Cu(InGa)Se~2~ film was peeled from it. The
peaks are identified in Table 17.

[]{#__RefHeading___Toc506348109 .anchor}Table 17 Identification of peaks
in Figure 49 by comparison to JCPDS card files for Mo, MoSe~2~, and
Cu(InGa)Se~2~.

  --------- -------- ------------ ------------ ----------------- ---------------------- -------------------
  **No.**   **2-θ\   **Counts**   **d (Å)\     **d (Å) (hkl) for                        
            °**                 measured**   identified phase                         
                                               and JCPDS file**                         

                                               **Mo(42-1120)**   **MoSe~2~(29-0914)**   **CIGS(35-1102)**

  1         13.64    293          6.486        --                6.46 (002)             --

  2         27.20    116          3.276        --                --                     3.312 (112)

  3         31.74    75           2.817        --                2.845 (100)            --

  4         38.40    55           2.342        --                --                     --

  5         40.56    4332         2.222        2.2247 (110)      --                     --

  6         44.96    13           2.014        --                --                     2.028 (220/204)

  7         56.10    20           1.638        --                1.643 (110)            --

  8         58.70    1480         1.571        1.5738 (200)      --                     --
  --------- -------- ------------ ------------ ----------------- ---------------------- -------------------

This technique was used to compare CuInSe~2~, Cu(InGa)Se~2~, and
Cu(InAl)Se~2~ films deposited on SLG and BSG substrates and also
correlated to adhesion. A number of films were peeled from the substrate
and XRD scans measured using a Cr tube for higher resolution. However,
it was found that they were not reproducible and changed with time. This
has been attributed to the reaction of the MoSe~2~ with the oxygen, and
probably H~2~O, present in the atmosphere. Measurement procedures have
been modified to minimize surface oxidation and experiments will be
repeated.

### 3.2.2 Sulfur diffusion in Cu(InGa)Se~2~

A Ph.D dissertation from the University of Delaware, Materials Science
Department titled "Surface Sulfurization of Copper Indium Diselenide for
Application in Solar Cells" was completed by Jochen Titus. The
dissertation abstract follows:

CuInSe~2~-based chalcopyrite films can be used as absorber layers in
thin-film solar cells. Some industrial manufacturing recipes of this
type of solar cell involve a separate sulfurization step after absorber
formation. The effects of this step on the absorber material, on its
electronic properties and on device performance are only poorly
understood. This project is part of a larger effort aimed at an
understanding of the mass transport, the changes in material properties
and the implications for device performance resulting from the
application of this technique. For this project, the objective was an
understanding of the mass transport taking place during the
sulfurization of CuInSe~2~ films, in particular the resulting
distribution of S in the film, which replaces some of the Se in the
chalcopyrite crystal lattice.

Samples were sulfurized in a flowing, H~2~S-containing atmosphere, under
conditions determined by annealing time and temperature and by gas
composition and flow. Experiments were done on films as well as on
single crystals. Samples were investigated with a range of analysis
tools, most notably Auger electron spectroscopy (AES) depth profiling,
X-ray diffraction (XRD), energy dispersive x-ray spectroscopy (EDS),
scanning electron microscopy (SEM) and atomic force microscopy (AFM).

Even subtle differences in sample-starting material were found to lead
to vastly different modes of S incorporation. Rapid S incorporation was
always correlated with a breakdown of the crystal lattice and with
excess Cu. It occurred in films containing excess Cu, as well as in
single crystals grown with the traveling heater method (THM) and with In
as solvent. In these crystal samples, the excess Cu became apparent only
after the sulfurization. A selenizing pre-anneal was effective in
preventing the lattice breakdown in these crystals during sulfur
reaction.

In single-phase film, as well as a vertical gradient freeze (VGF)-grown
crystal, a slow S incorporation mode was active. There, the S
distribution appeared to be the result of a fast chemical reaction at
the solid-gas interface, which saturated the sample surface with S, and
of subsequent solid-state interdiffusion of Se and S. In these samples
the crystal structure was conserved during sulfurization. It was shown,
that the S distribution in VGF crystals was the result of
one-dimensional diffusion, governed by a concentration-dependent binary
lattice diffusion coefficient d1, despite 4 elements (Cu, In, Se, S)
being present. The value of D1 was determined, as a function of S
concentration, with the help of the Bolzmann-Matano method from AES
depth profiles.

The recording, of a meaningful AES depth profile on a sulfurized film,
requires a film with a smooth front surface that must be conserved
during sulfurization. After preliminary experiments on rough films,
several attempts were made at preparing smooth films, since such films
were not readily available. Chemo-mechanical polishing appeared to the
most promising method, but it still needs to be developed further, for
experiments geared towards the determination of the values of both the
lattice and the grain boundary diffusion coefficient in these films. For
the evaluation of such experiments, a model of the interdiffusion in a
film was implemented in the software FEMLAB.

Na-containing single-phase films behaved very similar to Na-free
single-phase films. They also exhibited the slow S incorporation mode,
but in addition, they segregated CuInS~2~ and NaInS~2~ on their free
surface. The CuInS~2~ was distributed evenly across the sample surface,
but the NaInS~2~ segregated in a particular pattern.

In addition to the fundamental diffusion studies described in the
abstract above, a study of S diffusion into Cu(InGa)Se~2~ thin films as
a means to improve device performance was completed. The voltage of
Cu(InGa)Se~2~ solar cells can be increased by increasing the bandgap of
the absorber layer, although this gives a corresponding loss of
photocurrent. In order to reduce the recombination current and increase
the voltage without loss of photocurrent, inhomogeneous structures with
increased bandgap in the space charge region of the absorber layer could
potentially be utilized. Grading the space charge region of the absorber
layer with the addition of S should be preferable to Ga because more of
the bandgap increase results from a lowering of the valence band in the
case of S addition. Sulfur introduction into CuInSe~2~ and Cu(InGa)Se~2~
absorber layers by annealing these films in a H~2~S atmosphere has been
reported.[^46]^,^[^47]

In the present study, the post deposition surface sulfur reaction is
carried out in order to diffuse S into the space charge region of
homogeneous evaporated CIGS films. Polycrystalline Cu(InGa)Se~2~ thin
films were deposited by elemental evaporation on soda lime glass at
T~ss~=550°C and had Cu/(In+Ga)\~0.8-0.9 and Ga/(In+Ga)\~0.3. The sulfur
reactions were carried out in a flowing gas reactor at atmospheric
pressure with total flow rate 76 cm/min through the reaction chamber. To
control S diffusion and the S profile in the final film, Cu(InGa)Se~2~
samples were reacted in three different gas mixtures including hydrogen
sulfide, hydrogen selenide and oxygen diluted in argon.

As a baseline reaction procedure for sulfur reaction of Cu(InGa)Se~2~
films, the reaction was carried out in a 2% H~2~S-Ar mixture at 525°C
for various reaction times. A second process used the baseline reaction
conditions but with additional oxygen in the proportion O~2~/H~2~S=0.01.
Lastly, in order to slow down the sulfur reaction at the surface, the
reaction was performed in a mixed gas of hydrogen sulfide and hydrogen
selenide. This was done with the ratios H~2~S/(H~2~S+H~2~Se)=0.985, 0.98
and 0.97 with oxygen in the concentration O~2~/(H~2~S+H~2~Se)= 0.01 and
a total concentration of (H~2~S+H~2~Se)/Ar = 2%. All the reactions were
carried out for 120 min.

X-ray diffraction profiles of the (112) peaks before and after reaction
in H~2~S+O~2~ are shown in Figure 50. The as-deposited sample has a
narrow peak at 2θ = 27.04° the position expected for Cu(InGa)Se~2~
with Ga/(Ga+In) = 0.3. After sulfur reaction, the peak broadened due to
the incorporation of sulfur. A second peak appears at 2θ = 28.14°
suggesting the formation of a fully reacted Cu(InGa)S~2~ layer after the
reaction. The presence of the sulfide phase at the surface has been
verified using GIXRD. Comparable results were also obtained with
reaction in pure H~2~S using the baseline process as with reaction in
H~2~S + O~2~.

![](media/image64.pct){width="3.78125in" height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348200 .anchor}Figure 50 XRD spectra showing
the (112) peak before and after reaction in H~2~S + O~2.~

Reaction in H~2~S+H~2~Se+O~2~ resulted in broadening of the (112) peak
but not formation of a separate peak. This is shown in Figure 51 for
H~2~S/(H~2~S+H~2~Se)=0.98. This suggests the incorporation of sulfur to
form a mixed layer of Cu(InGa)(SeS)~2~ near the surface of the absorber.

![](media/image65.pct){width="3.78125in" height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348201 .anchor}Figure 51 XRD spectra showing
the (112) peak before and after reaction in H~2~S + H~2~Se + O~2.~

Auger electron spectroscopy (AES) depth profiles of S and Se are shown
in Figure 52 for the three reaction cases. After reactions in H~2~S or
H~2~S + O~2~ the S signal falls off sharply from the surface. Reaction
in H~2~S + H~2~Se + O~2~ leads to a much deeper incorporation of S and
gradients of S and Se. The AES depth profile results are consistent with
XRD results.

![](media/image66.pct){width="3.78125in" height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348202 .anchor}Figure 52 AES depth profiles of
Cu(InGa)Se~2~ films reacted in H~2~S, H~2~S + O~2~, and H~2~S + H~2~Se +
O~2~.

To further characterize the films, EDS measurements were done with 8kV
accelerating voltage so that only the top \~ 0.3 µm of the films is
probed. The Ga/(Ga+In) ratios were found to be 0.29, 0.22, and 0.04 for
the samples reacted in H~2~S, H~2~S + O~2~, and H~2~S + H~2~Se + O~2~
respectively. This indicates a significant decrease in relative Ga
content for the sample reacted in H~2~S + H~2~Se + O~2~. However, the
EDS analysis does not distinguish whether the Ga moves to the back or
moves out of the film. Figure 53 compares the AES depth profiles of the
relative Ga concentration for all the processes. In the as-deposited
Cu(InGa)Se~2~ film, the Ga concentration is uniform through the film.
For the sample reacted in H~2~S + H~2~Se + O~2~ there is depletion of Ga
from the surface and accumulation at the back.

![](media/image67.pct){width="3.78125in" height="3.1493055555555554in"}

[]{#__RefHeading___Toc506348203 .anchor}Figure 53 Ga/(In+Ga) ratios from
AES depth profiles of Cu(InGa)Se~2~ films reacted in H~2~S, H~2~S +
O~2~, and H~2~S + H~2~Se + O~2~.

In summary, post-deposition sulfur reactions of CIGS film have been
carried out in hydrogen sulfide with additional oxygen in the reaction
chamber and also in a mixture of hydrogen sulfide, hydrogen selenide and
oxygen. Oxygen incorporation promotes the decomposition of the hydride
gas and increases the subsequent sulfur incorporation into the film. A
fully converted sulfide layer is formed at the surface of the absorber
layer when reacted in hydrogen sulfide atmosphere, which can form a
barrier for current collection in device. With the mixed hydride gas a
graded Cu(InGa)(SeS)~2~ layer on the top of absorber layer was formed
and at the same time the Ga diffuses away from the surface. The low Ga
content after reaction means that the benefit of S incorporation near
the surface will not be obtained in devices. The bandgap near the
surface layer and in the space charge region of the device is not
controllably increased with S diffusion in this process so it does not
present a promising avenue for increased V~oc~.

## **3.3 Fundamental Interface and Material Characterization**

### 3.3.1 Cu(InGa)Se~2~ optical characterization

### 

We have previously determined the optical constants of polycrystalline
thin film Cu(InGa)Se~2~ alloys with Ga/(Ga+In) ratios from 0 to 1 using
variable angle spectroscopic ellipsometry (SE) over an energy range from
0.75 to 4.6eV.[^48] The films characterized in that work were all
nominally single phase Cu(InGa)Se~2~, with Cu/(In+Ga) ≈ 0.9 (\~24 at. %
Cu). Subsequently, we have focused on the effect of Cu stoichiometry in
Cu-poor materials on the optical properties. For the Cu-In-Se system,
the CuSe~2~-In~2~Se~3~ quasi-binary phase diagram determined by Gödecke
et. al. [^49] over the region of interest is shown in Figure 54. There
are single-phase regions containing the stoichiometric CuInSe~2~ α phase
and the δ~R~ phase, which includes the CuIn~3~Se~5~ composition. High
efficiency devices are made using films in the α or mixed phase α+δ~R~
regions. The objective in this study is to determine how the optical
functions vary, particularly in the α+δ~R~ region and across the very
wide single phase δ~R~ region, and whether the behavior in the mixed
phase region could be accurately fit as a mixture of the single phased
constants with the relative volume fractions determined by the fit.

Cu(InGa)Se~2~ thin films, deposited on soda lime glass/Mo substrates,
with Ga/(In+Ga) = 0.3 were used for this initial study. Their Cu
concentrations varied from 10 to 25 at .%, as shown in Figure 54,
corresponding to Cu/(In+Ga) from 0.3 to 0.93 and the compositions are
listed in Table 18. Sample preparation and the SE measurements were the
described elsewhere.Error: Reference source not found It is important to
note that the phase diagram in Figure 54 is only verified for the
Cu-In-Se system and there may be differences when Ga is included.

![](media/image68.pct){width="4.0in" height="3.3333333333333335in"}

[]{#__RefHeading___Toc506348204 .anchor}Figure 54 Cu~2~Se-In~2~Se
quasi-binary phase diagram taken from Reference Error: Reference source
not found. Triangles show the compositions of Cu(InGa)Se~2~ films used
in this study.

To describe the optical constants of a heterogeneous system a model is
needed and 2 cases are illustrated in Figure 55(a) shows a coated-sphere
microstructure. The dielectric response can be described by a Maxwell
Garnett effective medium approximation given by the relation:

$\frac{e - e_{a}}{e + 2e_{a}} = f_{b}\frac{e_{b} - e_{a}}{e_{b} + 2e_{a}}$
Equation 17

where ε~a~ and ε~b~ are, respectively, the dielectric function of two
phases a and b, ε is the effective dielectric function of the
heterogeneous medium, and f~b~ is the fraction of phase b. Figure 55(b),
shows a random mixture microstructure in which case the dielectric
response can be described by a Bruggeman effective medium approximation
given by the relation:

$0 = f_{a}\frac{e_{a} - e}{e_{a} + 2e} + f_{b}\frac{e_{b} - e}{e_{b} + 2e}$
Equation 18

Initially, the optical properties of the alloy films were determined
using multilayer optical modeling assuming the film is single phase.
Based on XRD measurements, this assumption is true for
Cu(In~0.7~Ga~0.3~)Se~2~ and Cu(In~0.7~Ga~0.3~)~3~Se~5~ compositions.
However, this assumption was later shifted and the mixed phase region
was analyzed with effective medium approximations using optical
constants from the end point compositions.

\(a\) ![](media/image69.png){width="2.348611111111111in"
height="1.4597222222222221in"}

\(b\) ![](media/image70.pct){width="2.2291666666666665in"
height="1.4395833333333334in"}

[]{#__RefHeading___Toc506348205 .anchor}Figure 55 Schematic of a
coated-sphere microstructure (a) and a random mixture of 2 phases (b).

Figures 56 and 57 show the index of refraction (n) and extinction
coefficient (k) obtained from the measurements after substrate and
surface roughness effects were removed using a multi-layer optical
model. Samples in the mixed phase region were analyzed with effective
medium approximations using optical constants from the end point
compositions. Two distinctive features are visible in these optical
spectra. First, the fundamental bandgaps are shifted to higher energies
with decreasing Cu concentration, which can be related to the decreasing
coupling between the Cu d- and Se p-orbitals, resulting in lowering of
the valance band maxima. Second, the critical point features at higher
energies become broader as the Cu concentration decreases. Broadening is
normally associated with degradation of the crystalline quality of the
material. Table 18 shows the volume fractions of the α and δ~R~ phases
using the Maxwell Garnet and Bruggeman EMA analyses. Similar results are
obtained with the different models for reasons that are not clear.

[]{#__RefHeading___Toc506348110 .anchor}Table 18 Compositions of
Cu-In-Ga-Se~2~ samples for SE analysis and the volume fraction of the α
and δ~R~ phases determined by the Maxwell Garnet and Bruggeman EMA
analyses. MSE is the mean square error in the analysis.

  ------------ --------- --------- --------- --------- ------------- ------------- ----------- ---------- --------- --------------- ---------- ---------
  **Sample**   **Cu**    **In**    **Ga**    **Se**    **Cu/**       **Ga/**       **Maxwell                        **Bruggeman**              
                                                                                   Garnett**                                                   

  **\#**       **(%)**   **(%)**   **(%)**   **(%)**   **(In+Ga)**   **(In+Ga)**   **α**       **δ~R~**   **MSE**   **α**           **δ~R~**   **MSE**

  33398        24.8      18.3      8.3       48.7      0.94          0.31          100         0                    100             0          

  33579        23.6      19.9      7.3       49.2      0.87          0.27          100         0          10        100             0          9.7

  33313        22.0      19.8      8.2       50.0      0.79          0.29          77          23         3.7       80              20         4.7

  33538        20.7      19.6      9.4       50.3      0.71          0.32          52          48         2.8       52              48         2.8

  33197        18.6      19.6      10.3      51.5      0.62          0.34          35          65         5.1       35              65         5.0

  33769        15.4      22.4      9.7       52.6      0.48          0.30          32          68         4.1       32              68         4.0

  33391        11.6      23.6      9.9       54.9      0.35          0.30          17          83         16.7      17              83         16.6

  33786        10.4      24.2      11.0      54.4      0.30          0.31                                           0               100        
  ------------ --------- --------- --------- --------- ------------- ------------- ----------- ---------- --------- --------------- ---------- ---------

![](media/image71.png){width="3.8583333333333334in"
height="3.1555555555555554in"}

[]{#__RefHeading___Toc506348206 .anchor}Figure 56 Index of refraction
spectra obtained by EMA analysis, of Cu-In-Ga-Se alloy films with
varying Cu concentration.

![](media/image72.png){width="3.8583333333333334in"
height="3.1555555555555554in"}

[]{#__RefHeading___Toc506348207 .anchor}Figure 57 Extinction coefficient
spectra obtained by EMA analysis, of Cu-In-Ga-Se alloy films with
varying Cu concentration.

The optical constants reported above depend on the Ga/(In+Ga) ratio as
well as the relative Cu concentration, so the determination of the
volume fractions of the different phases may be affected by small
variation in the Ga ratio. Hence a set of Cu-In-Se films was deposited
by elemental co-evaporation to enable characterization of films, in
which the optical properties are dependent only on the Cu ratio, x ≡
Cu/In. The composition of the films determined by EDS measurements are
listed in Table 19 and shown relative to the CuSe~2~-In~2~Se~3~
pseudo-binary phase diagramError: Reference source not found in Figure
58. The films were deposited on Mo-coated soda lime glass substrates and
then peeled as above. Ellipsometer measurements were again performed
immediately after peeling.

[]{#__RefHeading___Toc506348111 .anchor}Table 19 Composition of the
Cu-In-Se films used for the ellipsometry analysis.

  ------------ -------------- -------------- -------------- --------------
  **Sample**   **Cu**         **In**         **Se**         **x = Cu/In**

  33485.21     25.2           25.6           49.3           1.00

  33807.21     21.2           27.9           51.0           0.76

  33808.21     18.2           30.0           51.9           0.61

  33811.32     12.6           33.2           54.2           0.38

  33817.32     11.3           34.4           54.3           0.33
  ------------ -------------- -------------- -------------- --------------

![](media/image73.pct){width="3.8381944444444445in"
height="3.277083333333333in"}

[]{#__RefHeading___Toc506348208 .anchor}Figure 58 Cu~2~Se-In~2~Se~3~
pseudo-binary phase diagram taken from Reference Error: Reference source
not found. Triangles show the compositions of CuInSe~2~ films used in
this study.

The measured optical data, ε~1~ and ε~2~, was modeled using the method
described above. Optical constants for the single-phase α-CuInSe~2~ and
δ~R~-Cu~0.33~In~0.77~Se~2~ films are fit using oscillator levels in the
optical models. The optical data in the two-phase α+δ~R~ region was
modeled using a Bruggeman effective medium approximation. The optical
constants thus obtained, are shown in figures 59 and 60. These show a
decrease in dielectric response and broadening of critical point
features, with decreasing Cu ratio.

With both Cu-In-Se and Cu-In-Ga-Se alloy films an optical model has been
developed which allows the optical constants of mixed phase α+δ~R~
materials to be described by a mixture of the single phases of α and
δ~R~. Conversely, the composition of a mixed phase film could be
determined using the SE measurements. In both alloy systems there is an
increase in the energy of the latest order transitions, which determine
the bandgap. Also for each alloy system, there is a broadening of
critical point features with decreasing Cu, which indicates decreasing
crystalline quality.

![](media/image74.pct){width="5.129861111111111in"
height="3.245833333333333in"}

[]{#__RefHeading___Toc506348209 .anchor}Figure 59 Index of refraction
spectra for CuInSe~2~ films obtained from the oscillator and EMA models.

![](media/image75.pct){width="4.967361111111111in"
height="3.1506944444444445in"}

[]{#__RefHeading___Toc506348210 .anchor}Figure 60 Extinction coefficient
spectra for CuInSe~2~ films, obtained from the oscillator and EMA
models.

### 3.3.2 Emitter layer deposition and junction characterization

The effect of the Cu(InGa)Se~2~/emitter layer interface and band
alignment on devices is being studied by changing the alloy compositions
of the Cu(InGa)Se~2~ layers and comparing CdS, (CdZn)S or ZnS emitter
layers. By independently changing both the emitter and absorber bandgap,
it can be determined if the junction band alignment affects device
performance. With the wider bandgap in particular, it has been proposed
that the interface band alignment might limit V~oc~ [^50] and raising
the emitter layer conduction band could improve V~oc~. In addition,
different emitter deposition processes, which can affect the surface, or
interface chemistry, are being compared.

We have previously developed CBD processes for depositing the ZnS[^51]
and (CdZn)S[^52] layers. The (CdZn)S layer has bandgap 2.6 eV. In
addition, we have developed a process for CdS using chemical surface
deposition (CSD), which allows very high utilization of cadmium, 80-90%,
compared to utilization \< 10% in a typical CBD process.[^53] Earlier
results with CSD CdS in Cu(InGa)Se~2~ devices showed efficiencies
slightly lower than control devices with CBD CdS. In this work, devices
using Cu(InGa)Se~2~ cells with different bandgap absorber layers have
been prepared using CBD CdS, (CdZn)S, and ZnS and using CSD CdS and
(CdZn)S emitter layers.

Table 20 shows cell results using different CBD emitter layers deposited
on Cu(InGa)Se~2~ films with Ga/(In+Ga) = 0.3 and 0.6 corresponding to
bandgaps of 1.2 and 1.4 eV respectively. The results for the best cell,
out of 12 for each case, are summarized in the table. Cell results show
that on average, the (CdZn)S gives comparable performance to the CdS
emitter layer with both absorber layer bandgaps. The ZnS emitter layers
gave poor device efficiency in all cases and the process needs to be
re-evaluated.

[]{#__RefHeading___Toc506348112 .anchor}Table 20 Solar cell parameters
for Cu(InGa)Se~2~ with E~g~ = 1.2 or 1.4 eV using CdS, (CdZn)S and ZnS
emitter layers.

  ----------- --------- ------------- ------- ----------- -------------- ------- ------------- --------------
  **Sample\   **E~g~\   **emitter**   **η\    **V~oc~\    **J~sc~\       **FF\   **R~oc~\      **G~sc~\
  \#**        (ev)**                  (%)**   (Volts)**   (mA/cm^2^)**   (%)**   (Ω-cm^2^)**   (mS/cm^2^)**

  50392.23    1.2       CdS           15.9    0.627       33.3           76.0    1.3           1

  50392.33    1.2       (CdZn)S       15.7    0.620       32.7           77.2    1.3           2

  50392.31    1.2       ZnS           10.3    0.547       31.9           59.1    2.4           1

  50403.32    1.4       CdS           12.3    0.757       22.2           72.9    2.6           2

  50403.21    1.4       (CdZn)S       13.9    0.740       24.9           75.3    2.1           1

  50403.22    1.4       ZnS           4.7     0.613       18.7           40.5    10.8          9

  50405.32    1.4       CdS           12.8    0.765       22.8           73.8    2.3           1

  50405.21    1.4       (CdZn)S       13.0    0.739       23.8           73.6    2.3           1

  50405.12    1.4       ZnS           6.5     0.594       23.5           46.4    5.3           9
  ----------- --------- ------------- ------- ----------- -------------- ------- ------------- --------------

Additional experiments were done using (CdZn)S layers deposited by both
CBD and CSD processes to further determine whether there might be any
advantage to the (CdZn)S layers and to further evaluate the CSD process.
QE measurements are very useful to characterize the thickness and
bandgap of the (CdZn)S layers. The blue response in the QE curves for 4
samples from a single Cu(InGa)Se~2~ run are shown in Figure 61. The
curves shown were measured at OV under white light bias and were all
normalized to 1 at the maximum response which occurs at \~ 550 nm. E~g~
was determined from the wavelength shift at QE~norm~ ≈ 0.9. It is
assumed that the CdS has E~g~ = 2.42 eV. Then Figure 61 shows an E~g~
shift of 0.1 ± 0.03 eV so both the CBD and CSD (CdZn)S films have E~g~ =
2.52 eV. This corresponds to a composition with x ≡ Zn/(Cd+Zn) = 0.12
using a bowing parameter of 0.6. The target bandgap was 2.6 eV and the
Zn concentration will be adjusted in future experiments.

The CdS and (CdZn)S thicknesses were determined from the value of
QE~norm~ at λ = 400 nm assuming that the loss in the normalized QE is
due exclusively to CdS absorption and no light absorbed in the emitter
layer is collected. Then:

$\text{QE}_{\text{norm}}\left( \text{400}\text{nm} \right) = \frac{T}{1 - R}|_{\text{CdS}} = \text{exp}( - \alpha d)$.
Equation 19

This calculation used α = 1.45x10^5^ cm^-1^ for the CdS and 1.34x10^5^
cm^-1^ for the Cd~1‑x~Zn~x~S. Results are listed in Table 21. The
thicknesses for the CdS layers agree very closely to those on annealed
ITO control substrates.

![](media/image76.pct){width="3.1465277777777776in"
height="3.1465277777777776in"}

[]{#__RefHeading___Toc506348211 .anchor}Figure 61 Normalized short
wavelength QE used to determine thickness and E~g~ of CdS and
Cd~1‑x~Zn~x~S emitter layers deposited on Cu(InGa)Se~2~.

[]{#__RefHeading___Toc506348113 .anchor}Table 21 Thicknesses (d) of CdS
and Cd~1‑x~Zn~x~S determined from the normalized QE.

  ------------------------- --------------------- -------------- ----------------------
  **emitter layer**         **QE~norm~(400nm)**   **d(CdS)**     **d(Cd~1‑x~Zn~x~S)**

  CBD CdS                   0.47                  52             --

  CBD Cd~1‑x~Zn~x~S         0.35                  --             78

  CSD CdS                   0.33                  76             --

  CSD Cd~1‑x~Zn~x~S         0.51                  --             50
  ------------------------- --------------------- -------------- ----------------------

Device results comparing narrow 1.2 eV \[Ga/(In+Ga) ≈ 0.3\] and wide 1.5
eV \[Ga/(In+Ga) ≈ 0.75\] bandgap absorber layers with CdS and (CdZn)S
emitter layers deposited by CBD and CSD are summarized in Table 22 for
the best cell with each case. All results are after a 2 min air heat
treatment at 200°C which resulted in an increased V~oc~ and η for the
wider bandgap cells with little difference for the 1.2 eV cells. The
different emitter layers generally gave comparable J~sc~ and FF but the
CBD CdS always resulted in the highest V~oc.~ Thus there was no
advantage to the (CdZn)S emitter layers.

[]{#__RefHeading___Toc506348114 .anchor}Table 22 Solar cell parameters
for Cu(InGa)Se~2~ with E~g~ = 1.2 or 1.5 eV with CdS and (CdZn)S emitter
layers deposited by CBD or CSD.

  ------------------- ----------- ------------- -------- ----------- -------------- --------
  **Cu(InGa)Se~2~**   **Sample\   **CdS\        **η\     **V~oc~\    **J~sc~\       **FF\
                      \#**        type**        (%)**    (Volts)**   (mA/cm^2^)**   (%)**

  1.2 eV              33847.23    CBD CdS       13.6     0.609       30.3           73.5

                      33847.12    CSD CdS       12.4     0.575       29.9           72.0

                      33847.32    CBD CdZnS     11.9     0.557       30.0           71.1

                      33847.31    CSD CdZnS     10.8     0.558       27.7           70.1

  1.5 eV              33857.23    CBD CdS       9.2      0.782       18.3           64.4

                      33857.12    CSD CdS       8.8      0.700       18.3           68.4

                      33857.32    CBD CdZnS     7.1      0.595       18.4           65.1

                      33857.31    CSD CdZnS     7.5      0.641       18.1           64.5
  ------------------- ----------- ------------- -------- ----------- -------------- --------

A potentially critical difference between the CSD and CBD processes may
be due to etching or other chemical treatment of the Cu(InGa)Se~2~
surface which occurs in the first stage of CBD growth. In the CBD
process, the films are inserted into the bath containing thiourea and
ammonium hydroxide for 1 min prior to addition of CdSO~4,~ which
initiates the reaction to form CdS. Evidence of a surface etch from the
ammonia is provided by optical characterization of a Cu(InGa)Se~2~ film
which has been oxidized by exposure to humid air in the laboratory. In
this case, the optical parameters measured by spectroscopic ellipsometry
cannot be fit with a standard optical model.Error: Reference source not
found After a 1 min etch in 28 ml of 30% NH~4~OH diluted in 150 ml
H~2~O, identical to the concentrations in the CBD process, the
measurements fit the optical model developed for peeled Cu(InGa)Se~2~
with minimal exposure to air. This suggests that the ammonia etches an
oxide layer on the Cu(InGa)Se~2~. In the CSD process, the solutions
including the CdSO~4~ are applied directly to the hot Cu(InGa)Se~2~
surface and the reaction begins immediately.

Cu(InGa)Se~2~ from 2 different runs, with Ga/(In+Ga) ≈ 0.25, had CdS
deposited by baseline CBD, and CSD deposited at 65°C in 2 steps, i.e.
deposit, rinse and repeat deposition. The CdS was done with and without
a pre etch in ammonia solution as described above. During the CSD
process it was observed that the wetting of the etched Cu(InGa)Se~2~ was
different than the un-etched films. Specifically, a low contact angle
was observed on the un-etched Cu(InGa)Se~2~ suggesting a higher surface
energy. With the NH~3~-etched Cu(InGa)Se~2~ a high contact angle was
observed similar to deposition on glass, suggesting lower surface
energy.

The results for the best cell on each piece are summarized in Table 23
and the uniformity among 6 cells on each piece was good. In each case,
the efficiency with CSD CdS, with or without the pre-etch, is comparable
within statistical uncertainty, to the CBD samples. However, the V~oc~
is \~ 30 mV lower with the CSD process. JV and QE curves show no
qualitative differences in behavior but more detailed J-V analysis is
being completed. In addition, some of the samples were given a 2 min air
heat treatment at 200°C but there was no significant difference.

While the difference in V~oc~ is small it has been observed in several
experiments and shows that the CdS deposition process has a complicated
effect on device performance. Further experiments will evaluate other
surface Cu(InGa)Se~2~ treatments prior to CSD deposition to help
understand the relationship between the surface chemistry and electronic
properties.

[]{#__RefHeading___Toc506348115 .anchor}Table 23 Effects of NH~3~
etching Cu(InGa)Se~2~ prior to CSD CdS.

  ----------- ------------ --------------- --------- ----------- -------------- ---------
  **CdS\      **Sample\    **condition**   **η\      **V~oc~\    **J~sc~\       **FF\
  Process**   \#**                         (%)**     (Volts)**   (mA/cm^2^)**   (%)**

  CBD\        33846.23     as-dep          12.7      0.616       28.6           72.2
  control                                                                       

              33847.23     as-dep          12.9      0.605       29.2           73.1

              33847.23     HT\*            13.6      0.609       30.3           73.5

  CSD         33846.12     as-dep          12.3      0.584       30.4           69.4

              33847.12     as-dep          11.9      0.580       28.7           71.6

              33847.12     HT\*            12.4      0.575       29.9           72.0

  NH~3~ / CSD 33846.13     as-dep          12.6      0.582       30.0           72.5

              33847.13     as-dep          12.3      0.584       28.3           74.4

              33847.13     HT\*            12.9      0.571       30.9           73.1
  ----------- ------------ --------------- --------- ----------- -------------- ---------

HT\* - heat-treated for 2 min at 200°C in air.

# 4 Si-Based Solar Cells

## **4.1 Fabrication of Large Grain Si Films**

### 4.1.1 Post-deposition annealing experiments 

### 

Post-deposition annealing experiments of Al/Si bi-layers were carried
out to compare crystallization and layer inversion processes in standard
metal induced crystallization and in-situ metal induced crystallization.
The use of evaporated Si films also allows comparison between films
deposited in an H free environment and HWCVD films. First, 0.5 mm thick
films of Al were deposited onto 7059 glass substrates by electron beam
evaporation. The films were left in ambient air for a period of 24 hr to
allow formation of an oxide layer. Although the role of the oxide layer
is not well understood, literature reports indicate its presence is
required to achieve layer inversion and a continuous polycrystalline Si
film. Si films with 0.6 mm thickness were then deposited by electron
beam evaporation and HWCVD. XRD measurements showed no evidence of
crystallization of the Si films deposited by evaporation prior to
annealing (Figure 62). On the other hand, XRD and Raman data indicated
complete crystallization of the films deposited by HWCVD (Figure 63).
Figure 64 is an optical micrograph through the glass substrate of the
sample deposited by HWCVD. In this picture, the bright regions represent
Al while the dark regions are Si grains at the glass interface. These
reflectivity observations have been corroborated by Raman spectroscopy.
Figure 64 clearly shows that a continuous Si film was not obtained at
the end of deposition. The films deposited by HWCVD were thus added to
the annealing set to find if further layer inversion could be achieved.
Annealing of all films was performed in a tube furnace for 90 min at
500°C under 100 sccm of Ar. For the samples with Si films deposited by
evaporation, optical microscopy showed polycrystalline Si nucleation
extending to the glass interface (Figure 65). However for the annealing
time and temperature used, there was very little lateral grain growth.
Consistent with this observation, XRD analysis showed only partial
crystallization of these films (Figure 62). For samples with Si films
deposited by HWCVD, no increase in grain coalescence or crystalline
fraction was observed. All of these samples were further annealed at the
same conditions, for an additional 5 hrs. However, no change in
properties was observed after annealing.

![](media/image77.jpeg){width="3.698611111111111in"
height="3.321527777777778in"}

[]{#__RefHeading___Toc506348212 .anchor}Figure 62 XRD pattern of
evaporated Si film on Al before and after annealing.

![](media/image78.jpeg){width="3.670138888888889in"
height="3.347916666666667in"}

[]{#__RefHeading___Toc506348213 .anchor}Figure 63 XRD pattern of HWCVD
Si film on Al before and after post deposition annealing.

![](media/image79.jpeg){width="3.563888888888889in"
height="2.654166666666667in"}

[]{#__RefHeading___Toc506348214 .anchor}Figure 64 Optical micrograph
through the glass substrate of Si film deposited by HWCVD.

![](media/image80.jpeg){width="3.265972222222222in"
height="2.5097222222222224in"}

[]{#__RefHeading___Toc506348215 .anchor}Figure 65 Optical micrograph
through the glass substrate of Si film deposited by evaporation after 90
min anneal at 500 °C.

### 4.1.2 In-situ aluminum induced crystallization experiments

Analysis of HWCVD Si films deposited on Al coated glass substrates below
the Si-Al eutectic temperature, has revealed a layered microstructure.
the lower layer near the glass interface is composed of micron size
islands of polycrystalline silicon. Atop this layer, a mixture of
nanocrystalline Si and Al remains (Figure 66a). This mixture is also
observed in the spaces between the polycrystalline Si islands (Figure
66b). The existence of this layered structure is evidence of layer
inversion and is consistent with standard metal induced crystallization
results obtained by other groups. However, the discontinuity of the
underlying polycrystalline Si structure limits its usefulness for device
fabrication. Consequently, experiments were carried out to study the
effect of process parameters on the continuity of the polycrystalline Si
layer.

![](media/image81.jpeg){width="3.061023622047244in"
height="2.3255905511811026in"}![](media/image82.jpeg){width="3.072834645669291in"
height="2.322834645669291in"}

[]{#__RefHeading___Toc506348216 .anchor}Figure 66 Nanocrystalline Si
layer atop a) and between b) polycrystalline Si islands.

Deposition parameters and film properties are summarized in Table 24.
All Si films were deposited onto 0.5 μm Al layers at 430 °C and 1 μm/hr
growth rate. Since the presence of an oxide on the Al layers is a
critical factor in achieving layer inversion, the time the Al samples
were exposed to air was also recorded. Figure 67 shows a typical film as
seen by optical microscopy through the glass side. Raman spectroscopy
through the glass substrate (Figure 68) confirms that the dark regions
are composed polycrystalline Si while the bright regions are a mixture
of nanocrystalline Si and Al. After etching the overlaying Si-Al
mixture, a similar pattern of polycrystalline Si islands is observed
from the front surface. This combined with the similarity of the Raman
profile of these islands with that of c-Si suggests that the islands are
individual Si grains.

Images similar to Figure 67 were taken at equal magnification for all
samples and analyzed using image processing software to determine the
fractional coverage of the polycrystalline Si layer and the average
grain size. From Table 24 as the Si/Al ratio increases (samples
234-236), the fractional coverage also increases. Since the supply of Si
increases for these samples, it is reasonable to expect that the
fractional coverage would increase. However, diffusion of Si through the
Al layers should be rate limiting compared to the delivery rate of Si to
the Al layer surface. Consequently, the increase in fractional coverage
could be an effect of the longer annealing times (deposition times)
considering that in all cases the supply of Si is sufficient to replace
the Al layer. To verify this, another experiment was performed to
replicate sample 235. However, the new film (sample 241) was subjected
to a post-deposition anneal at 430°C under vacuum for 2 hr. The image of
the resulting film is shown in Figure 69. The fractional coverage
increases from 60 to 90%. This indicates that the excess Si does not
crystallize as it accumulates on the Al layer surface during deposition
and is available for further lateral growth of the polycrystalline Si
layer. Further experiments are needed to determine the minimum
post-deposition annealing time to achieve adequately continuous
polycrystalline films.

The average grain size as shown in Table 24 appears to decrease with
increasing Si/Al ratio (samples 234-236). However, the results for
samples 241 and 237 clearly show this is not the case. On the other
hand, the average grain size correlates well with the time the Al layers
remained exposed to air prior to Si film deposition.

![](media/image83.jpeg){width="3.666535433070866in"
height="2.8019685039370077in"}

[]{#__RefHeading___Toc506348217 .anchor}Figure 67 Optical micrograph
through the glass substrate where the dark features are Si grains.

![](media/image84.png){width="3.089583333333333in" height="3.15in"}

[]{#__RefHeading___Toc506348218 .anchor}Figure 68 Raman profiles of
polycrystalline Si grains and nanocrystalline Si/Al mixture compared to
that of c-Si.

![](media/image85.jpeg){width="3.7188976377952754in"
height="2.739763779527559in"}

[]{#__RefHeading___Toc506348219 .anchor}Figure 69 Optical micrograph of
deposited film following annealing for 2 hr at 430°C.

The role of aluminum oxides on the layer inversion process and grain
size of the polycrystalline Si layer, although not well understood, has
been acknowledged in previous work of standard metal induced
crystallization. It is hypothesized that the oxide layer acts as a
nucleation barrier leading to a lower density of near neighbor nuclei;
thus more lateral space for grains to grow before they are impinged upon
by the closest grain. However, it is expected that the oxide layer is
self-limiting and would reach equilibrium after a brief period of time
of exposure to air. The present results suggest that the oxide layer
continues to change long after deposition. Interestingly, the time of
exposure to air does not influence the resulting fractional coverage as
sample 237 shows.

![](media/image86.png){width="2.6666666666666665in"
height="1.9333333333333333in"}
![](media/image87.png){width="2.6666666666666665in"
height="1.9333333333333333in"}

![](media/image88.png){width="2.6666666666666665in"
height="1.9333333333333333in"}
![](media/image89.png){width="2.6666666666666665in"
height="1.9333333333333333in"}

[]{#__RefHeading___Toc506348220 .anchor}Figure 70 Optical micrographs
through the glass substrate of deposited Si films on a) 100 nm, b) 250
nm, c) 500 nm and d) 1000 nm Al.

[]{#__RefHeading___Toc506348116 .anchor}Table 24 Process parameters and
film properties.

+----------+-----+----------+--------------+---------+---------------+
| >        | ### | > **De   | > **Time of  | >       | > **Average   |
| **Sample | ### | position | > Al layer** |  **Frac | > grain       |
| > \#**   |  Si | > time** | >            | tionalc | > size**      |
|          | /Al | >        | > **exposure | overage | >             |
|          |     | >        | > to air     | > (%)** | > **(μm)**    |
|          |     | **(hr)** | > (days)**   |         |               |
+----------+-----+----------+--------------+---------+---------------+
| HW 234   | 1   | 0.5      | 80           | 50      | 10            |
+----------+-----+----------+--------------+---------+---------------+
| HW 235   | 2   | 1        | 81           | 60      | 10            |
+----------+-----+----------+--------------+---------+---------------+
| HW 236   | 4   | 2        | 1            | 70      | 4             |
+----------+-----+----------+--------------+---------+---------------+
| HW 241   | 2   | 1        | 50           | 90      | 7             |
+----------+-----+----------+--------------+---------+---------------+
| HW 237   | 1   | 0.5      | 2            | 50      | 4             |
+----------+-----+----------+--------------+---------+---------------+

In order to consider the effect of growth rate, as well, another set of
Si films (HW 245) with thickness of 2-3 μm was deposited at 5 μm/hr on
Al layers 100, 250, 500 and 1000 nm thick. The deposition temperature
was 430°C. Figures 70a, 70d are optical micrographs taken through the
glass substrate of the Si films deposited on the four thickness values
of Al. The dark areas in the micrographs represent polycrystalline Si
grains or film where grains have coalesced. Contrary to previous results
at 1μm/hr growth rate, the fractional coverage in this case increases
with decreasing Si/Al ratio. The central area of the film on 500 nm Al,
appears to have the densest coverage approaching 100%. The edge areas of
this film had a fractional coverage resembling that of Figure 66b.
However, after the 8 hr post deposition anneal at 500°C, the edge area
had a fractional coverage similar to that in Figure 70c. Such large
gains in fractional coverage could not be achieved with films deposited
at 1μm/hr. Perhaps this could be explained by the nature of the Si layer
deposited on the Al. At higher growth rates, it is conceivable that the
Si films deposited were amorphous as a higher silane flow rate and lower
residence time were used. Amorphous Si defuses readily into Al and leads
to complete MIC of the Si deposited. On the other hand, the
microcrystalline; which would result from lower deposition growth rates
does not interact with Al at 430°C leading to incomplete MIC of the Si
deposited. The films also had a significant number of cracks, which are
visible in some of the micrographs. It is not clear at this point if
this is the result of the higher deposition rates or of stress in the Al
layers.

These continuous polycrystalline layers (Figure 70c) appear as promising
candidates for device fabrication. Future work will include electrical
characterization and use of the layers as templates for epitaxial growth
of large grain i-layers. Additional results on this section were
included in a paper titled "In-situ aluminum induced crystallization of
silicon thin films on glass substrates using HW-CVD" at the Spring 2004
MRS meeting in San Francisco.

### 4.1.3 Role of Al film interface

In previous experiments, it has been observed that the average grain
size of in-situ aluminum-induced crystallized (AIC) Si films correlates
well with the time the Al layers remained exposed to air prior to Si
film deposition. Figure 71 shows optical microscopy images of in-situ
aluminum-induced crystallized samples taken through the glass
substrates. It has been reported that the aluminum oxide layer is
required for the layer exchange process although it is not required for
the crystallization process itself. It is reasonable to claim that the
oxide layer acts as a physical barrier forcing the growth of Si grains
laterally when they reach the thickness of the initial aluminum layer.
It also acts a as diffusion barrier for Si and Al atoms leading to a
lower nucleation and thus large Si grains. There has been little
attention paid to the characterization of the oxide layer, other than
some work from Nast at UNSW.[^54] Since the growth of the oxide depends
on time between Al and Si depositions, Al surface treatments and Al
sample storage, it represents a potentially uncontrolled variable in
much of the AIC work.

The oxide surface was monitored over a period of 12 days using
spectroscopic ellipsometry measurements. Samples of e-beam deposited Al
on glass were stored in air. Although formation of the native oxide
layer is thought to be a self-limiting process, ellipsometry results in
Figure 72 reveal changes in the optical properties of the oxide layer
indicating that its structure may be changing long after the Al film
deposition.

![](media/image90.jpeg){width="5.942361111111111in"
height="1.4777777777777779in"}

[]{#__RefHeading___Toc506348221 .anchor}Figure 71 Optical microscopy
pictures of partially crystallized Si samples taken through the glass
substrate: (a) with Al layer exposed to air for 24 hr before Si
deposition, (b) with 48 hrs, and (c) with 80 days.

> ![](media/image91.jpeg){width="3.7152777777777777in"
> height="3.1305555555555555in"}

[]{#__RefHeading___Toc506348222 .anchor}Figure 72 Ellipsometry
measurements showing a progressive change in the optical parameter ε~2~
of 100 nm thick aluminum layers even after 12 days. Inset shows the
decrease in the peak over the range 1.45 to 1.7 eV.

In order to understand the effect of the initial aluminum layer on the
crystallization and layer exchange process, aluminum layers with
different thicknesses were deposited by electron-beam deposition on 7059
glass substrates. XPS and Auger analysis were then performed on aluminum
films. As expected, XPS depth profiles showed an oxygen rich surface on
the aluminum layers as shown in Figure 73. The plot is normalized with
respect to the Al-O peak, which is located at 75.6 eV. As the sputtering
depth increases the intensity of the bulk aluminum peak, which is
located at 73.0 eV, also increases. However, a detailed analysis of the
spectra shows a shift for the Al-O peak to a lower binding energy as the
sputtering depth increases indicating the presence of more than one
oxidation state.

> ![](media/image92.jpeg){width="6.0in" height="5.041732283464567in"}

[]{#__RefHeading___Toc506348223 .anchor}Figure 73 XPS depth profile for
Al2p for 100 nm thick Al film on glass. The binding energies for Al-Al
and Al-O are 73.0 eV and 75.6 eV, respectively. Data is normalized with
respect to the Al-O peak.

The O1s peak also undergoes a similar shift as seen in the XPS depth
profile shown in Figure 74. The shift to higher energies indicates the
presence of a higher oxidation state, perhaps Al(OH)~3~. This suggests
that the oxide layer is a bi-layer structure; composed of Al(OH)~3~ and
on Al~2~O~3~. Auger analysis appears to agree with the XPS results.
Figure 75 shows an Auger depth profile of Al films with thicknesses of
100, 250, 500 and 1000 nm. The Al/O ratio on the sample surface changes
from 1 to approximately 50 after several sputtering cycles. However this
ratio does not show a sharp increase indicating the absence of discrete
layers, as also suggested by XPS measurements.

![](media/image93.jpeg){width="4.375196850393701in"
height="4.104330708661418in"}

[]{#__RefHeading___Toc506348224 .anchor}Figure 74 XPS depth profile for
O1s showing a shift to a lower energy as the sputtering depth increases.
The energies observed are in good agreement with the data from the
literature considering the data shown here is not corrected nor
normalized.

Previously, we have reported that such high temperature Si deposition on
Al was problematic since the Al films appear to become discontinuous.
Work in progress is directed at optimizing the growth of Si films at
high temperatures, above the Si-Al eutectic. Recently, it was found that
Al films heated to 600°C in vacuum remained continuous. Thus the problem
may be related to de-wetting of the liquid phase formed at the eutectic.
Buffer layers between the glass and Al will be studied to see if they
can maintain wetting on the glass hence give continuous Al films with Si
film deposition at T\>600°C. Additionally, a series of Si films have
been deposited by HWCVD, on c-Si substrates under conditions which may
lead to epitaxial growth. Samples are being prepared for TEM
measurements. These experiments will provide conditions to grow large
grain Si films with controlled, doping on AIC Si templates.

![](media/image94.jpeg){width="5.159722222222222in"
height="2.9895833333333335in"}

[]{#__RefHeading___Toc506348225 .anchor}Figure 75 Auger depth profiles
showing the change in composition of the aluminum oxide layers.

## **4.2 Solar Cell Development**

### 4.2.1 Impurities in a-Si layers

Previously it was reported that a-Si devices fabricated in the single
chamber PECVD reactor had lower fill factors compared to devices
fabricated several years ago. Earlier devices routinely achieved FF=70%
while recently the FF has been limited to 50-60%. QE analysis of the
recent devices indicates that the lower fill factors are caused by poor
hole collection in the i-layer, most likely arising from donor
impurities such as P or O. This could impact our ability to make
contacts and intrinsic buffer layers for various device structures
required to evaluate the HWCVD films and devices. Several a-Si i-layers
and devices were sent to NREL for SIMS characterization to compare
levels of impurities before and after the decrease in fill factor. Table
25 lists the samples sent for analysis and their characteristics. Since
we are concerned about residual contamination, the previous run in the
reactor is also of importance and is noted in the table. Our standard
procedure is to deposit a 15 min a-SiC "burying" layer between devices
to eliminate dopant cross-contamination, specifically P from the
previous n-layer being incorporated into the next devices p-layer. Run
41032 was intentionally deposited after a n-layer deposition but without
a burying layer.

The concentrations of impurities such as B, Cl, F, Al, Ba, Na, K, Ca,
Sn, Fe and Cr, were below the background levels observed in a c-Si
wafer. The N profiles are shown in Figure 76. Concentrations of the
atmospheric contaminants N and C (not shown) in the i-layer of all
samples were comparable in magnitude and relatively uniform in
distribution. The apparent thickness difference between more recent
samples and sample 4767 is due to the higher texture of the TCO in
sample 4767 which leads to an effective broadening of the profiles at
the SnO~2~/a-Si interface. While C could have other sources, N is only
available from an atmospheric leak. Figure 76 suggests that a leak to
atmosphere is not responsible for the changes we have observed,
consistent with the low leak-up rate.

![](media/image95.jpeg){width="4.5in" height="3.2694444444444444in"}

[]{#__RefHeading___Toc506348226 .anchor}Figure 76 Nitrogen depth
profiles.

[]{#__RefHeading___Toc506348117 .anchor}**Table 25 a-Si i-layers and
devices characterized by SIMS.**

+--------+---------------------+--------+----------------------------+
| **     | **Structure**       | **Date | **Comments**               |
| Sample |                     | Depos  |                            |
| \#**   |                     | ited** |                            |
+--------+---------------------+--------+----------------------------+
| 4767   | 7059 glass/0.3 m   | 02     | a-Si device deposited      |
|        | SnO~2~:F/           | /10/98 | after std. a-SiC burying   |
|        |                     |        | run. Device had high FF    |
|        | 0.03 m B-doped     |        |                            |
|        | a-Si/               |        |                            |
|        |                     |        |                            |
|        | 0.01 m a-SiC/      |        |                            |
|        |                     |        |                            |
|        | 0.5 m undoped      |        |                            |
|        | a-Si/               |        |                            |
|        |                     |        |                            |
|        | 0.03 m P-doped     |        |                            |
|        | a-Si                |        |                            |
+--------+---------------------+--------+----------------------------+
| 41028  | SL glass/0.3 m     | 09     | a-Si device deposited      |
|        | SnO~2~:F/           | /03/03 | after std. a-SiC burying   |
|        |                     |        | run. Device had low FF     |
|        | 0.03 m B-doped     |        |                            |
|        | a-Si/               |        |                            |
|        |                     |        |                            |
|        | 0.01 m a-SiC/      |        |                            |
|        |                     |        |                            |
|        | 0.5 m undoped      |        |                            |
|        | a-Si/               |        |                            |
|        |                     |        |                            |
|        | 0.03 m P-doped     |        |                            |
|        | a-Si                |        |                            |
+--------+---------------------+--------+----------------------------+
| 41032  | 7059 glass/0.2 m   | 09     | a-Si i-layer deposited     |
|        | Cr/                 | /22/03 | without a prior burying    |
|        |                     |        | run and after an n-layer   |
|        | 0.5 m undoped a-Si |        | deposition                 |
+--------+---------------------+--------+----------------------------+
| 41034  | 7059 glass/0.2 m   | 09     | a-Si i-layer deposited     |
|        | Cr/                 | /30/03 | after a std. a-SiC burying |
|        |                     |        | run                        |
|        | 0.5 m undoped a-Si |        |                            |
+--------+---------------------+--------+----------------------------+
| 41035  | 7059 glass/0.2 m   | 10     | a-SiC p-layer, a-SiC       |
|        | Cr/                 | /03/03 | buffer and a-Si i-layer    |
|        |                     |        | deposited after std. a-SiC |
|        | 0.03 m B-doped     |        | burying run and a-Si       |
|        | a-Si/               |        | i-layer deposition         |
|        |                     |        |                            |
|        | 0.01 m a-SiC/      |        |                            |
|        |                     |        |                            |
|        | 0.5 m undoped a-Si |        |                            |
+--------+---------------------+--------+----------------------------+

![](media/image96.jpeg){width="4.1305555555555555in"
height="3.0027777777777778in"}

[]{#__RefHeading___Toc506348227 .anchor}Figure 77 Oxygen depth profiles
for all samples plus P profile for 41032 (dashed line).

Figure 77 shows the oxygen depth profiles. More recent samples, except
for 41032, contain oxygen concentrations a factor of two greater than
the device with high FF (4767). In addition, all of these samples have a
gradient, which decreases from the glass interface, as if the source of
contamination is being depleted as the layer was deposited. This
gradient is absent in sample 4767. For sample 41032, which was deposited
after an n-layer deposition and without a prior a-SiC
contaminant-burying run, the difference in O concentrations with those
in sample 4767 is an order of magnitude. The gradient in the profile is
significant. The absence of similar gradients in the C and N
concentrations for all the samples indicates that the O does not
originate from atmospheric contamination, e.g., from a leak. For all
samples, except 41032, the P concentrations were below the c-Si
background level. The P profile in sample 41032 shown by the dashed blue
curve in figure 77 also decreases from the glass interface as the P is
being depleted. The trend and increased magnitude of the O and P
profiles in sample 41032 is similar, suggesting a common source. These
results suggest that the P contamination could be higher now but below
the SIMS limit, but can be tracked by the increased O. Re-optimizing the
burying layer to reduce the P was investigated next.

Several runs were made with varying burying layer and bake-out
procedures using a same device recipe, which gave 9-10% devices several
years ago. Since then, dopant gas dilution decreased from 2% to 0.2%. As
a result, flows and times had changed. Due to miscommunication, the
devices made after the new burying layer had a thicker p-layer than
intended, i.e. 75 sec instead of the standard 20 sec. This serendipitous
error actually gave some important insight into the contamination issue.

Table 26 shows the solar cell parameters for the recent device 41074 (75
sec p-layer) to a device 41030 deposited September 2003 (20 sec
p-layer), which also had the typical low FF (55%). They are nearly
identical. The more recent cell has slightly lower J~sc~ (as expected
due to thicker p-layer) and V~oc~ and slightly higher FF.

[]{#__RefHeading___Toc506348118 .anchor}Table 26 J-V parameters for a-Si
pin solar cells having different p-layer times. QE is shown in Fig. 78a
and 78b.

  ----------------------- ----------- ----------- -------- --------- ----------- -----------
  **cell**                **V~oc~**   **J~sc~**   **FF**   **Eff**   **R~oc~**   **G~sc~**

  41030 (20 sec p-layer)  0.856       13.5        56       6.5       11.6        2

  41074 (75 sec p-layer)  0.834       12.0        59       5.9       11.3        2
  ----------------------- ----------- ----------- -------- --------- ----------- -----------

Figure 78a shows the QE at 0V, -1V, and +0.5V on the older device 41030.
The QE at 400 is similar for all biases, but at longer wavelengths the
QE decreases substantially with increasing forward bias. Figure 78b
shows the QE for the new device 41074. The QE at 400 nm is lower than in
Figure 78a, consistent with the much thicker p-layer. But the curves
have much greater voltage dependence at short wavelengths than at longer
wavelengths, opposite to Figure 78a. Figure 79 shows the QE bias ratios
explicitly. Red lines are the 20 sec p-layer and black lines are the new
cell with 75 sec p-layer. Note how the bias dependence is nearly
symmetrical and opposite from each other. A common signature of the low
FF we had been seeing for over 2 years was large bias dependence to the
QE at long wavelengths. This is a common indication of poor hole
collection as due to a donor impurity or low field. Capacitance
measurements and reasonably good V~oc~ somewhat ruled out the low field,
leaving low lifetime as the likely cause. The new device with longer B
flow has the opposite limitation. This indicates that the longer p-layer
contributed significant B dopant, which was not removed by our standard
p-i interface pump/purge cycle and a-SiC buffer layer. This step had
been optimized with the standard p-layer. Apparently, it was
insufficient to remove all the B from the chamber after the longer
p-layer. But the crucial point is that there is no long wavelength bias
dependence.

In general, this experiment confirms that devices in a single chamber
system are sensitive to dopant cross-contamination. Our device process
had evolved by establishing a careful balance between impurities in our
single chamber reactor. Tipping that balance slightly towards donor or
acceptor impurities results in a contaminated i-layer and low FF. It is
encouraging that the FF is slightly higher now. This could mean we have
actually eliminated the donor-like impurity (probably O) or that it is
being partially compensated. It is even possible that we could actively
compensate the donor defects by adjusting the B flow or time.

![](media/image97.pct){width="2.9097222222222223in"
height="3.0034722222222223in"}
![](media/image98.pct){width="2.9097222222222223in"
height="3.0034722222222223in"}

[]{#__RefHeading___Toc506348228 .anchor}Figure 78 (a) and 78 (b). QE
curves at --1V, 0V, and +0.5V bias for devices with 20 sec (left) and 75
sec (right) p-layer.

![](media/image99.pct){width="3.63125in" height="3.7527777777777778in"}

[]{#__RefHeading___Toc506348229 .anchor}Figure 79 QE ratios for devices
with 20 and 75 sec p-layers.

### 4.2.2 Boron doped HWCVD Si films 

The deposition of boron doped HWCVD Si films was investigated as part of
our effort to develop polycrystalline thin film Si devices. The films
were deposited at 500 m Torr and 450°C with B~2~H~6~/SiH~4~ flow ratios
varying from 10^-5^ to 10^-3^. For all dopant gas concentrations, the
resulting dark conductivity of the films was 10^-6^ S/cm. In order to
understand whether the invariance in conductivity was the result of the
deposition conditions or film structure, additional doped films were
deposited. Deposition conditions and film properties for this sample set
are given in Table 27, which also includes one sample (HW217-21) from
the previous experimental set. Comparison of samples HW223-11 and
HW217-21 clearly indicates that the deposition pressure does not play a
role in affecting the film conductivity. For sub-micron 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 samples HW226-11
and HW217-21. 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.

Additional films were deposited in the thickness range of 30-50 nm in
order to develop doped films for window layers. Preliminary results show
that all films have conductivities of 10^-2^ S/cm for B~2~H~6~/SiH~4~
flow ratios of 0, 1x10^-2^ and 4x10^-2.^ The invariance in conductivity
between undoped and doped films at these dopant gas concentrations
probably reflects a dominating effect of the film surface on the
conductivity values. Other groups have reported this effect for films in
this thickness range. For future work, doped films in the range of 200
nm thickness will be deposited to study the bulk properties of films
adequate for window layers.

[]{#__RefHeading___Toc506348119 .anchor}Table 27 Deposition conditions
and properties of boron doped HWCVD Si films.

+-----------+---------+-------------------+-------------+-------------+
| **Sample  | **Thic  | **Deposition**    | **Con       | **A         |
| \#**      | kness** |                   | ductivity** | ctivation** |
+-----------+---------+-------------------+-------------+-------------+
|           | *       | **Conditions**    | **(S/cm)**  | **Energy    |
|           | *(mm)** |                   |             | (eV)\***    |
+-----------+---------+-------------------+-------------+-------------+
|           |         |                   |             |             |
+-----------+---------+-------------------+-------------+-------------+
| HW217-21  | 0.5     | P =200 mTorr      | Dark=1.1E-6 | Not         |
|           |         | B~2~H~6~/         |             | Measured    |
|           |         |                   |             |             |
|           |         | SiH~4~ = 1.0x10   |             |             |
|           |         | ^-3^              |             |             |
|           |         |                   |             |             |
|           |         | SiH~4~ = 4 sccm   |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Tf = 1850 ^o^C    | Photo=same  |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Ts = 450 ^o^C     |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Gr = 1.0 μm/hr    |             |             |
+-----------+---------+-------------------+-------------+-------------+
| HW223-11  | 1.5     | P =500 mTorr      | Dark=1.0E-6 | 0.088       |
|           |         | B~2~H~6~/         |             |             |
|           |         |                   |             |             |
|           |         | SiH~4~ = 1.0x10   |             |             |
|           |         | ^-3^              |             |             |
|           |         |                   |             |             |
|           |         | SiH~4~ = 4 sccm   |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Tf = 1850 ^o^C    | Photo=same  |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Ts = 450 ^o^C     |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Gr = 3.0 μm/hr    |             |             |
+-----------+---------+-------------------+-------------+-------------+
| HW226-11  | 0.8     | P =200 mTorr      | Dark=1.5E-3 | 0.041       |
|           |         | B~2~H~6~/         |             |             |
|           |         |                   |             |             |
|           |         | SiH~4~ = 1.0x10   |             |             |
|           |         | ^-2^              |             |             |
|           |         |                   |             |             |
|           |         | SiH~4~ = 3 sccm   |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Tf = 1850 ^o^C    | Photo=same  |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Ts = 450 ^o^C     |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Gr = 1.2 μm/hr    |             |             |
+-----------+---------+-------------------+-------------+-------------+
| HW227-11  | 2.5     | P =200 mTorr      | Dark=1.1E-1 | 0.040       |
|           |         | B~2~H~6~/         |             |             |
|           |         |                   |             |             |
|           |         | SiH~4~ = 1.0x10   |             |             |
|           |         | ^-2^              |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | SiH~4~ = 3 sccm   |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Tf = 1850 ^o^C    | Photo=same  |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Ts = 450 ^o^C     |             |             |
+-----------+---------+-------------------+-------------+-------------+
|           |         | Gr = 1.2 μm/hr    |             |             |
+-----------+---------+-------------------+-------------+-------------+

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

# In-line process diagnostics

## **5.1 Introduction**

The in-line process diagnostic effort is intended to develop the tools
needed for process control and quantitative models to link sensor output
to process variables to material properties. During this year, the
contact wetting angle was explored in detail. In this section, the
contact wetting angle measurement is presented as a simple diagnostic
tool that may be suitable in industrial applications for monitoring
changes in thin film solar cell processing conditions. While applicable
to any thin film solar cell, it has been developed here primarily with
application to CdTe/CdS devices. The previous annual report presented
successful application to Cu(InGa)Se~2~/CdS devices.

The cross-section of a thin film CdTe/CdS device is given in Figure 80
and depicts each thin film layer in sequential order of deposition
as: 1) indium tin oxide (ITO), 2) CdS, and 3) CdTe. Methods currently
utilized to improve device performance require that each thin film layer
be subjected to post-deposition treatments referred to as processing
steps. Processing is broadly defined as any technique used to improve
device performance and in this report will be considered to be heat
treatments to ITO in atmospheres of air and argon, and CdCl~2~ heat
treatments followed by etching of CdTe in a bromine-methanol solution
(Br~2~:CH~3~OH).

![](media/image100.png)

[]{#__RefHeading___Toc506348230 .anchor}Figure 80 Schematic of the
cross-section of a CdTe solar cell detailing thin film layers in
sequential order of deposition.[^55]

In devices, post-deposition processing results in various beneficial
changes that derive from structural and chemical changes at the film
surface and within the bulk. Many of these effects have been well
documented for CdTe.[^56]^,^[^57] For example, thermal treatments with
exposure to CdCl~2~ are known to alter structural properties of CdTe by
changing crystalline grain orientation (primary recrystallization) and
initiating grain growth (secondary recrystallization). The extent of
these changes depends on the prior thermal history of the film and film
thickness, which control the as-deposited orientation and grain
size.[^58] For example our recent work with vapor transport (VT)
deposited films having \>2 μm grains has shown that CdCl~2~ treatment
does not measurably change film morphology yet is still required to
obtain optimal cell performance.[^59] For physical vapor deposited (PVD)
films initially consisting of small crystallites (\<1 μm) such as those
used in this study it has been shown that significant grain growth and
changes in grain orientation from (111) to random occur. Additionally,
the CdTe surface chemistry is affected by reaction with the ambient
species, CdCl~2~ and O~2~, resulting in formation of oxide species and
stoichiometric deviations:[^60]

*CdTe(s) + O~2~(g)  CdO(s) + TeO~2~(s)* Equation 20

*CdTe(s) + CdCl~2~(s/g) + O~2~(g)  2CdO(s) + TeCl~2~(g)* Equation 21

Reaction 20 may contribute to CdTe doping by shifting the surface
stoichiometry, by formation of cadmium vacancies due to the greater
thermodynamic driving force for Cd oxidation (-54 kcal/mol) compared to
that of Te (-15 kcal/mol). At the typical processing temperature of
400ºC, substantial redistribution of these vacancies into the bulk via
diffusion could facilitate acceptor complex formation by reaction with
impurity species such as Cl, O, and Cu.^58,^[^61] Electronically, this
results in greater p-type conductivity desired for low series resistance
and higher open circuit voltage (V~oc~).[^62] For cell processing,
surface treatment is often required to remove oxides formed by CdCl~2~
treatments, as they have insulating electrical effects. Many treatments
have been developed, and a common treatment consists of etching CdTe in
a solution of Br~2~:CH~3~OH, which preferentially removes Cd, leaving a
Te layer.[^63]^,^[^64] Removal of oxides and formation of the Te layer
are critical to obtaining a low resistance back contact. Less is known
regarding the specific effects of processing treatments on the surfaces
of ITO and CdS thin films, though it has been documented that resistance
to interdiffusion and chemical robustness are improved for crystallized
layers in ITO and CdS thin films.[^65]

This work investigates the effects of processing treatments on thin film
surfaces in order to further identify processing-induced changes which
can be used for diagnostics in commercial-scale fabrication of CdTe thin
film modules. Analytical methods employed at the laboratory scale are
typically expensive, sophisticated, and time-consuming techniques; are
inappropriate in a manufacturing environment and are not easily
integrated.[^66] Therefore, it is advantageous to seek less costly and
efficient diagnostic tools that offer quantitative information to gauge
the effectiveness of post-deposition treatments by the changes they
impart to the exposed film surface.

The terminating surface of a solid is known to contain excess free
energy when compared to the bulk, which arises in part from greater
molecular disorder at the termination of the bulk lattice. This property
is referred to as the surface energy of a solid and is experienced by
liquids in the form of surface tension. Surface energies consist of two
components that manifest themselves as polar energies, resulting from
permanent and induced dipole moments as well as hydrogen bonding, and
dispersive energies that arise from vibrational interactions between
atoms.[^67] The sum of these individual energy components represents the
overall surface energy for liquids and solids. The dependence of solid
surface energies to surface properties makes it a suitable candidate for
use as a tool in investigating the effects of post-deposition
processing.

The interaction of liquids and solids of different surface energies are
known to give rise to the formation of characteristic angles known as
wetting angles at a liquid/solid interface. This phenomenon was first
described by Young in 1804 when he presented an equation that quantifies
the energy of a solid surface by measuring the angle that the
cross-section of a liquid droplet of known surface tension makes when
contacting a solid substrate.[^68] This technique for measuring surface
energy is commonly referred to as contact wetting analysis. Chasse and
Ross have used this method to observe oxidation on silicon surfaces by
monitoring decreases in wetting angles[^69] and Lockhande et al. have
reported success with contact wetting as an empirical method for
evaluating CuInS~2~ thin films.[^70] Additionally, preliminary
observations in our laboratory consistently show substantial changes in
the hydrophilic/hydrophobic characteristics of CdTe films at different
stages of processing. Contact wetting angle analysis will therefore be
employed as a diagnostic tool to monitor changes in the surface energies
of thin films in a CdTe/CdS solar cell resulting from processing during
device fabrication and correlate them to alterations in surface
structure and chemistry as determined by GIXRD and XPS analyses.

## **5.2 Methodology**

### 5.2.1 Surface energy analysis 

Young's equation given below:^68^

*E~SV~ = E~SL~ + E~LV~ cos θ* Equation 22

relates the energy of a solid in air (*E~SV~*) to the energy of a liquid
in air (*E~LV~*) and at a solid liquid interface (*E~SL~*) by measuring
the wetting angle (*θ*), but does not take into account the existence of
polar ($E_{i}^{p}$) and dispersive ($E_{i}^{d}$) surface energies; where
the subscript *i* denotes either a liquid (1, 2) or a solid (s) surface.
Figure 81 shows a magnified CCD image of the cross-section of a liquid
droplet applied to a CdTe thin film and illustrates the wetting angles
(*θ~L~* and *θ~R~*) created by the interaction of the solid and liquid
surfaces.

Equation 23 gives the relationship between the overall surface energy of
a liquid or solid (*E~i~*) to the magnitude of ($E_{i}^{p}$) and
($E_{i}^{d}$):

$E_{i} = E_{i}^{p} + E_{i}^{d}$ Equation 23

This work employed the harmonic mean method to calculate
$E_{s}^{p}$and$E_{s}^{d}$. The resulting equations are given below:[^71]

$(1 + \text{cos}\theta_{1})E_{1} = 4(\frac{E_{1}^{d}E_{s}^{d}}{E_{1}^{d} + E_{s}^{d}} + \frac{E_{1}^{p}E_{s}^{p}}{E_{1}^{p} + E_{s}^{p}})$
Equation 24

$(1 + \text{cos}\theta_{2})E_{2} = 4(\frac{E_{2}^{d}E_{s}^{d}}{E_{2}^{d} + E_{s}^{d}} + \frac{E_{2}^{p}E_{s}^{p}}{E_{2}^{p} + E_{s}^{p}})$
Equation 25

![](media/image101.png)

[]{#__RefHeading___Toc506348231 .anchor}Figure 81 CCD image of the
cross-section of a water droplet deposited onto CdTe with left and right
contact angles represented by θ~L~ and θ~R~, respectively.

Two liquids of known polar and dispersive surface energies can be used
in conjunction with experimentally measured wetting angles (θ*~i~*) to
solve Equations 24, and 25 simultaneously to obtain
$E_{s}^{p}$and$E_{s}^{d}$. The components of surface energies for
liquids are easily measured and have been well documented for many
fluids. Table 28 gives the polar fraction (χ~p~) and total surface
energies (*E~1,2~*) for liquids used in this study.

[]{#__RefHeading___Toc506348120 .anchor}Table 28 Properties of liquids
used for contact wetting analysis, where T~b~ is the boiling
temperature, χ~p~ is the polar fraction and E~S~ is the total surface
energy at room temperature.[^72]

+---------------+---------------+------------------+------------------+
| **Liquid**    | **T~b~        | **χ~p~**         | **E~S~**         |
|               | (^o^C)**      |                  |                  |
|               |               |                  | **(mN/m)**       |
+---------------+---------------+------------------+------------------+
| Water         | 100           | 0.70             | 73.0             |
+---------------+---------------+------------------+------------------+
| Formamide     | 111           | 0.32             | 58.2             |
+---------------+---------------+------------------+------------------+

### 5.2.2 Contact wetting angle measurements

A droplet of liquid listed in Table 28 was deposited onto the surface of
a thin film that had been cleaned with isopropanol and dried with N~2~.
Liquids were deposited with a simple bulb dropper placed at a fixed
height approximately 2 cm above the sample. Magnified digital images of
the cross-section of droplets were collected after a minimum of 10 sec,
as suggested by Lokhande et al., to allow for complete interaction
between liquid and solid surfaces.^70^ Measurements were conducted in an
open atmosphere and temperature range of 25^o^C-28^o^C. To minimize
statistical variation, two to six wetting angle measurements were
conducted per sample with each fluid. The average of the left and right
angles of the cross-section of each droplet was taken to be the wetting
angle of that droplet (Figure 81). The characteristic wetting angle used
to calculate surface energies was the collective average of wetting
angles for each thin film sample, and the error bar for each measurement
was \~ ± 2.5^o^.

The edge-detect feature in the image analysis toolbox available in
Mathworks, Inc. software package Matlab^^ was employed to profile the
edges of droplets. The software package contains several different
detection methods that identify contrast differences in images to create
edge profiles. This work selected the Canny method, as it was the most
likely to detect weakly contrasted edges. Standard Procedures
Pixelport^^ shareware was then used in conjunction with edge profile
images to determine wetting angles.

### 5.2.3 Film preparation and characterization

ITO was radio frequency sputtered onto glass in an argon atmosphere of 5
Torr at room temperature using a sintered target with a tin oxide weight
fraction of 5%. The growth rate was 15 Å/min and film thickness was
approximately 2500 Å. ITO films were stored under ambient conditions.
Contact wetting angle and GIXRD were used to analyze changes in the
surfaces of ITO thin films resulting from heat treatments in atmospheres
of argon and air at a temperature of 500^o^C for 10 min.

CdTe films were deposited from 6N purity CdTe powder by physical vapor
deposition (PVD) at a temperature 325^o^C and a growth rate of
approximately 0.1 μm/min to obtain a film thickness of approximately 3
μm. Contact wetting was used in conjunction with GIXRD and XPS to detect
changes on the surfaces of PVD (111) CdTe thin films associated with
post-deposition treatments. Surface energies were measured at different
times with respect to removal from the deposition chamber. Selected
samples were stored for approximately 2.5 months under the following
conditions: vacuum desiccator (pH~2~O \<0.1 Torr), desiccator (pH~2~O= 5
Torr) and 1.5 months in saturated air (pH~2~O= 24 Torr). Partial
pressures of water vapor in air were calculated using the Antoine
equation and the percent humidity of the storage environment. Surface
energies of single crystal (111) Cd and Te faces that had been stored
under ambient conditions were compared to all samples.

The relative rate and short-term effects of native oxidation on CdTe
surface energies were monitored by immediately measuring the magnitude
of contact wetting upon removal from the evacuated PVD chamber and
subsequent to chemical etching. Changes were measured periodically over
the course of one hr. Exposure to CdCl~2~ was accomplished by subjecting
CdTe films to a solution of CdCl~2~ and methanol. Heat treatments were
conducted in air ambient at 400^o^C for 20 min.

Digital images of wetting angles were obtained at a magnification of 64X
with a SAC Imaging model 7 CCD camera that had been fit into the ocular
of a Gaertner light microscope. GIXRD measurements for all samples were
conducted with a Rigaku DMax 2200 diffractometer with CuK-α radiation in
asymmetric parallel-beam mode and an incident beam angle of
approximately 1^o^.[^73] XPS analyses were completed for CdTe films
using a Perkin Elmer PHI 560 ESA/SAM system equipped with an Al anode
and set to a power of 200W.

## **5.3 Results and Discussion**

### 5.3.1 GIXRD patterns and surface energies for as-deposited and heat-treated thin films of In~2~O~3~:SnO (ITO) 

Figure 82 shows the GIXRD patterns before and after thermal anneal for
2500 Å thick ITO sputtered onto soda lime glass. The low intensity of
defined peaks in the as-deposited GIXRD patterns indicates amorphous or
microcrystalline structure prior to processing. Heat treatments to ITO
in atmospheres of either air or argon for 10 min at 500^o^C sharpened
peaks signifying crystallization and are indexed as the sesquioxide
In~2~O~3~ (ICDD# 6-416). The sole cause of crystallization is thermal
activation as it occurred irrespective of atmospheric conditions. Table
29 compares surface energy components obtained for different processing
conditions. For atmospheres containing air and argon, the polar surface
energy component (E^p^) was increased by 21.5 and 22.3 mN/m,
respectively, when compared to the as-deposited sample. The dispersive
surface energy (E^d^) component was raised slightly by 2.1 and 3.2 mN/m
for air and argon treatments, respectively. Energy changes may be
attributed to higher surface polarity associated with the alignment and
organization of bonds from previously amorphous states.

![](media/image102.png){width="5.173611111111111in"
height="2.734027777777778in"}

[]{#__RefHeading___Toc506348232 .anchor}Figure 82 GIXRD patterns for
as-deposited and treated ITO. Treatments included air and argon anneal
at 500^°^C for 10 min. The sampling depth (1/e basis) for Cu K~ α~
X-rays in ITO at 1^o^ incidence is \~ 1000 A.

[]{#__RefHeading___Toc506348121 .anchor}Table 29 Surface energies for
as-deposited (As-Dep) and heat-treated (HT) ITO thin films. Energies are
listed as: dispersive energy (E^d^), polar energy (E^p^) and total
surface energy (E~S~). (∆E^d^) and (∆E^p^) signify changes in dispersive
and polar energies from as-deposited conditions, respectively.

+-------+---------+---------+----------+---------+---------+---------+
| *     | **Trea  | *       | **E^p^** | *       | **      | **      |
| *Thin | tment** | *E^d^** |          | *E~S~** | ∆E^d^** | ∆E^p^** |
| f     |         |         | **       |         |         |         |
| ilm** |         | **(     | (mN/m)** | **(     | **(     | **(     |
|       |         | mN/m)** |          | mN/m)** | mN/m)** | mN/m)** |
+-------+---------+---------+----------+---------+---------+---------+
| ITO   | As-Dep  | 19.8    | 8.3      | 28.1    | n/a     | n/a     |
+-------+---------+---------+----------+---------+---------+---------+
| ITO   | HT Air  | 21.9    | 29.8     | 51.7    | 2.1     | 21.5    |
+-------+---------+---------+----------+---------+---------+---------+
| ITO   | HT Ar   | 23.0    | 30.6     | 53.6    | 3.2     | 22.3    |
+-------+---------+---------+----------+---------+---------+---------+

### 5.3.2 Native oxidation of CdTe and its effect on surface energies 

As-deposited PVD CdTe films are (111) oriented and exhibit small grain
size \~ \< 1 μm. These properties do not appreciably change over time,
however, the surface wetting characteristics of the film continue to
evolve when exposed to room air. Figure 83 shows the rate at which the
wetting angle of deionized water changes over the course of an hr for a
CdTe thin film immediately upon exposure to room air from the evacuated
deposition chamber. Also presented in Figure 83 are the effects on
wetting angle subsequent to chemical etching with a bromine-methanol
solution. Figure 83 shows that the film tested upon removal from the
deposition chamber nearly completely wet the entire surface of the
sample (θ \~1^o^). Rapid increases in this value were observed over the
first hr. Bromine-methanol etched films, however, experienced differing
effects with respect to the magnitude of wetting angles, which increased
rapidly during the first 20 min, then subsequently decreased implying
that etching may not produce surfaces similar to that of as-deposited
CdTe. The rate of change in wetting angles was observed to be most
significant within the first hr of exposure to the ambient room air for
both as-deposited and etched CdTe implying that the majority of change
to (111) oriented CdTe occurred rapidly (\< 1hr). The results after
several hr on these samples are similar to results obtained for samples
measured after prolonged storage times.

![](media/image103.png){width="5.625in" height="4.770833333333333in"}

[]{#__RefHeading___Toc506348233 .anchor}Figure 83 Wetting angle of
deionized water vs. time for PVD (111) CdTe during the first hr of
exposure to the atmosphere from the evacuated deposition chamber (solid
circles) and after chemical etching of room air stored CdTe for 5 sec in
a solution of Br~2~:CH~3~OH (open squares).

CdTeO~3~ (ICDD \# 22-129) has been reported as the predominant native
oxide of CdTe on single crystals[^74] and changes in the magnitudes of
wetting angles can therefore be attributed to the formation of this
oxide on the surface of as-deposited CdTe films. For etched CdTe, on
which a Te layer is produced the predominant oxide is TeO~2~. The
results for etched CdTe can thus be interpreted dynamically in light of
previous work,^63^ in which a thin amorphous Te layer is generated which
begins to crystallize over a twenty-min period and is subsequently
oxidized to TeO~2~.

### 5.3.3 Effects of storage environments prior to processing on CdTe GIXRD patterns, XPS surface concentrations and surface energies.

### 

To determine the specific effects of native oxidation on surface energy
components over an extended time period, additional CdTe samples were
stored for over two months under vacuum desiccator (pH­~2~O \<0.1Torr)
and desiccator (pH~2~O= 5 Torr) conditions. Additionally, CdTe
thin-films were stored \~ 1.5 months in water saturated air (pH~2~O= 24
Torr) and single crystal (111) Cd and Te face samples were stored in air
ambient. The peak at 23.7^o^ given by GIXRD data in Figures 84 and 85
shows films to be (111) in orientation, and other peaks in the range of
5^o^-40^o^ reveal the presence of the native oxide CdTeO~3~. XPS data in
Table 30 report oxygen surface concentrations that increase with storage
humidities. Surface energies for as-deposited (111) CdTe and (111)
single crystal Cd and Te face samples are compared to storage conditions
in Table 31. Surfaces are shown to be energetically similar for all
samples, which is likely due to the presence of native oxides regardless
of prior storage.

![](media/image104.png){width="5.089583333333334in"
height="2.734027777777778in"}

[]{#__RefHeading___Toc506348234 .anchor}Figure 84 GIXRD patterns for PVD
CdTe stored under vacuum desiccator (pH~2~O \< 0.1 Torr) conditions for
\~2.5 months prior to processing. Processing treatments included CdCl~2~
treatment in air at 400^o^C for 20 min, followed by etching for 5 sec in
a solution of Br~2~:CH~3~OH. The sampling depth for Cu K~ α~ X-rays in
CdTe at 1^o^ incidence is \~ 1160 A.

![](media/image105.png){width="5.089583333333334in"
height="2.734027777777778in"}

[]{#__RefHeading___Toc506348235 .anchor}Figure 85 GIXRD patterns for PVD
CdTe stored under desiccator (pH~2~O= 5 Torr) conditions for \~2.5
months prior to processing. Processing treatments included CdCl~2~
treatment in air at 400^o^C for 20 min followed by etching for 5 sec in
a solution of Br~2~:CH~3~OH. The sampling depth for Cu K~ α~ X-rays in
CdTe at 1^o^ incidence is \~ 1160 A.

### 5.3.4 Effects of storage and processing on CdTe GIXRD patterns, XPS surface concentrations, and surface energies 

The effects of CdCl~2~ treatment and etching on CdTe films stored at
different humidity were analyzed by GIXRD, XPS and surface energy
analysis using contact wetting methods. The storage humidity ranged from
\<0.1 Torr for vacuum dessicator storage to 24 Torr for water-saturated
air. The films were treated with CdCl~2~ in dry air ambient at 400^o^C
for 20 min and subsequently etched for 5 sec in a solution of
Br~2~:CH~3~OH. Heat treatments with CdCl~2~ induced crystallite growth
and oxide formation (Figures 84-86). The formation of particular oxide
phases was found to be dependent on storage conditions prior to
treatment. For samples stored in "dry" air environments (pH~2~O \<0.1
Torr and pH~2~O= 5 Torr), GIXRD data in Figures 84 and 85 indicate the
presence of CdO (ICDD# 5-640) by the reflection at 33.5^o^ and
additional peaks representing CdTeO~3~ after processing. XPS
measurements of oxygen concentration on samples after storage and
CdCl~2~ treatment were 25 to 35% and were reduced significantly by
etching (Table III), suggesting a high degree of surface oxidation. This
is further supported by GIXRD data in Figures 84 and 85, which show
diminished CdTeO~3~ reflections and disappearance of CdO peaks. In
addition, the XPS data show a decrease in the Cd/Te ratio emphasizing
the effectiveness of this processing technique in selective removal of
Cd from the lattice. For samples stored in "wet" air environments
(pH~2~O = 24 Torr), the results differ from dry stored samples. Figure
86 shows GIXRD patterns for CdTe thin-films stored in water-saturated
air at 25^o^C (pH~2~O= 24Torr). The peaks at 30.5^o^ and 50.5^o^ shown
in Figure 86 exhibit the formation of CdTe~2~O~5~ (ICDD# 20-1404)
subsequent to CdCl~2~ heat treatments. The broad peak occurring at

\~ 30.5^o^ is indexed as two reflections, due to CdTe~2~O~5~ at 30.0^o^
and due to ITO in the substrate, at 30.6^o^. Although the XPS
measurements in Table 31 show a significant decrease in surface oxide
concentration after etching with Br~2~:CH~3~OH, the GIXRD patterns in
Figure 86 show that CdTe~2~O~5~ remains predominant. This indicates that
the oxide penetrated the bulk of the film, showing the strong effect of
storage humidity on the CdTe surface composition obtained during
processing.

![](media/image106.png){width="5.094444444444444in"
height="2.759027777777778in"}

[]{#__RefHeading___Toc506348236 .anchor}Figure 86 GIXRD patterns for PVD
CdTe stored in saturated air at 25^o^C (pH~2~O= 24 Torr) for \~1.5
months prior to processing. Processing treatments included CdCl~2~ heat
treatment in air at 400^o^C for 20 min followed by etching for 5 sec in
a solution of Br~2~:CH~3~OH. The sampling depth for Cu K~α~ X-rays in
CdTe at 1^o^ incidence is \~ 1160 A.

The effects of processing on the surface energy components of CdTe thin
films are shown in Table 31. The CdCl~2~ treatment primarily affects
E^P^ leaving E^d^ largely unchanged. Surface oxidation resulting from
CdCl~2~ treatment increased E^P^ by 9.8, 8.7, and 5.9 mN/m for samples
previously stored in vacuum, dry air, and saturated air, respectively.
Etching in Br~2~:CH~3~OH reduced E^P^ by 4.4, 16.8, and 9.6 mN/m due to
the removal of oxides from the film surfaces that had been stored under
vacuum, dry air, and saturated air, respectively, prior to treatments.
Although the wetting angle measurements are capable of revealing gross
changes in processing-induced oxidation, they are insufficient to
differentiate between different oxide phases. Overall, contact wetting
was found to be sensitive to oxide formation during CdCl~2~ treatment
and oxide removal by etching. After CdCl~2~ treatment, surface oxidation
makes the film more hydrophilic as shown in Figure 87a for a water drop
on CdCl~2~ treated CdTe. After etching, the surface becomes hydrophobic
as shown in Figure 87b. The sensitivity of water contact wetting is due
to its high polar energy component, making it a useful diagnostic tool
for monitoring the effects of CdTe processing.

The presence of CdTe~2~O~5~ found by GIXRD measurements has implications
for device performance, since formation of CdTe~2~O~5~
disproportionately removes Te from the lattice, thereby limiting the
ability to form Cd vacancies needed for p-type conduction. In light of
these findings and previous work;^59^ it is proposed that oxidation of
CdTe proceeds via *at least* three reactions that are dependent on
storage conditions prior to and during processing. The reactions are as
follows:

\(1\) *Native oxidation in dry air: CdTe + 3/2O~2~  CdO + TeO~2~ 
CdTeO~3~*

\(2\) *Native oxidation in humid air: CdTe + 5/2O~2~  CdO + 2TeO~2~ 
CdTe~2~O~5~*

> \(3\) *Reaction with CdCl~2~ in dry air: CdTe + CdCl~2~ + O~2~  2CdO
> +TeCl~2~*

The results above suggest that for all-dry storage and CdCl~2~
processing, reactions 1 and 3 predominate, while the addition of water
vapor to the ambient enhances reaction 2. Further work must be done to
investigate the specific effects of storage environments on device
operation and conversion efficiency.

[]{#__RefHeading___Toc506348122 .anchor}Table 30 XPS atomic percent
surface concentrations for PVD CdTe thin films stored under various
conditions.

+--------+--------------+-----+------+------+------+-----+----+----+
| > **S  | > *          | > * | >    | > *  | >    | >   | >  | >  |
| torage | *Treatment** | *C~ |  **C | *C~C | **C~ | **C |  * | ** |
| >      |              | TeI | ~TeI | d~** | O~** | ~Cl | *C | C~ |
| pH~2~O |              | ~** | I~** |      |      | ~** | ~C | O/ |
| > (T   |              |     |      |      |      |     | d~ | ~C |
| orr)** |              |     |      |      |      |     | /C | ~C |
|        |              |     |      |      |      |     | ~T | d~ |
|        |              |     |      |      |      |     | e~ | +C |
|        |              |     |      |      |      |     | ** | ~T |
|        |              |     |      |      |      |     |    | e~ |
|        |              |     |      |      |      |     |    | ** |
+--------+--------------+-----+------+------+------+-----+----+----+
| \<0.1  | None         | 9   | 1    | 2    | 2    | N/a | 1. | 0. |
|        |              | .2% | 1.8% | 1.3% | 7.3% |     | 01 | 64 |
+--------+--------------+-----+------+------+------+-----+----+----+
| \<0.1  | CdCl~2~      | 7   | 8.3% | 1    | 3    | 3   | 1. | 1. |
|        |              | .8% |      | 6.2% | 5.4% | .5% | 01 | 09 |
+--------+--------------+-----+------+------+------+-----+----+----+
| \<0.1  | CH~3~OH:     | 0   | 2    | 2    | 2    | 0   | 0. | 0. |
|        | Br~2~ Etch   | .0% | 4.8% | 2.4% | 3.2% | .0% | 91 | 49 |
+--------+--------------+-----+------+------+------+-----+----+----+
|        |              |     |      |      |      |     |    |    |
+--------+--------------+-----+------+------+------+-----+----+----+
| 5      | None         | 11  | 8.2% | 1    | 2    | N/a | 0. | 0. |
|        |              | .4% |      | 5.3% | 9.9% |     | 78 | 86 |
+--------+--------------+-----+------+------+------+-----+----+----+
| 5      | CdCl~2~      | 6   | 1    | 1    | 3    | 5   | 1. | 0. |
|        |              | .9% | 0.1% | 7.7% | 1.4% | .0% | 04 | 91 |
+--------+--------------+-----+------+------+------+-----+----+----+
| 5      | CH~3~OH:     | 0   | 2    | 2    | 1    | 0   | 0. | 0. |
|        | Br~2~ Etch   | .0% | 5.0% | 2.3% | 9.3% | .0% | 89 | 41 |
+--------+--------------+-----+------+------+------+-----+----+----+
|        |              |     |      |      |      |     |    |    |
+--------+--------------+-----+------+------+------+-----+----+----+
| 24     | None         | 9   | 2.3% | 1    | 3    | N/a | 1. | 1. |
|        |              | .1% |      | 6.8% | 6.2% |     | 48 | 28 |
+--------+--------------+-----+------+------+------+-----+----+----+
| 24     | CdCl~2~      | 7   | 7.7% | 1    | 3    | 10  | 1. | 0. |
|        |              | .2% |      | 9.6% | 1.7% | .0% | 32 | 92 |
+--------+--------------+-----+------+------+------+-----+----+----+
| 24     | CH~3~OH:     | 0   | 2    | 2    | 1    | 0   | 0. | 0. |
|        | Br~2~ Etch   | .0% | 7.4% | 5.3% | 6.9% | .0% | 93 | 32 |
+--------+--------------+-----+------+------+------+-----+----+----+

[]{#__RefHeading___Toc506348123 .anchor}Table 31 Surface energies for
as-deposited and treated CdTe thin films stored under various
conditions. Energies are listed as: dispersive (E^d^), polar (E^p^),
total surface energy (E~S~), change in dispersive subsequent to each
treatment (∆E^d^), and change in polar subsequent to each treatment
(∆E^p^).

+------+---------+--------+-------+-------+-------+-------+-------+
| > ** | > **St  | > *    | > **E | > **E | > **E | >     | >     |
| Thin | orage** | *Treat | ^d^** | ^p^** | ~S~** |  **∆E |  **∆E |
| > fi |         | ment** | >     | >     | >     | ^d^** | ^p^** |
| lm** |         | >      | >     | >     | >     | >     | >     |
|      |         | > **(m | **(mN | **(mN | **(mN | >     | >     |
|      |         | N/m)** | /m)** | /m)** | /m)** | **(mN | **(mN |
|      |         |        |       |       |       | /m)** | /m)** |
+------+---------+--------+-------+-------+-------+-------+-------+
| Cd   | ambient | As-Rec | 23.6  | 38.4  | 62    | n/a   | n/a   |
+------+---------+--------+-------+-------+-------+-------+-------+
| Te   | ambient | As-Rec | 20.2  | 40.4  | 60.6  | n/a   | n/a   |
+------+---------+--------+-------+-------+-------+-------+-------+
|      |         |        |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | \<      | As-Dep | 26.5  | 35.2  | 61.7  | n/a   | n/a   |
|      | 0.1Torr |        |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | \<      | C      | 25.5  | 45    | 70.5  | -1.0  | 9.8   |
|      | 0.1Torr | dCl~2~ |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | \<      | Etch   | 25.1  | 40.6  | 65.7  | -0.4  | -4.4  |
|      | 0.1Torr |        |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
|      |         |        |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | 5 Torr  | As-Dep | 25.6  | 38.7  | 64.3  | n/a   | n/a   |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | 5 Torr  | C      | 23.6  | 47.4  | 71    | -2.0  | 8.7   |
|      |         | dCl~2~ |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | 5 Torr  | Etch   | 27.7  | 30.6  | 58.3  | 4.1   | -16.8 |
+------+---------+--------+-------+-------+-------+-------+-------+
|      |         |        |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | 24 Torr | As-Dep | 23    | 39.9  | 62.9  | n/a   | n/a   |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | 24 Torr | C      | 25.4  | 45.8  | 71.2  | 2.4   | 5.9   |
|      |         | dCl~2~ |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+
| CdTe | 24 Torr | Etch   | 25.7  | 36.2  | 61.9  | 0.3   | -9.6  |
+------+---------+--------+-------+-------+-------+-------+-------+
|      |         |        |       |       |       |       |       |
+------+---------+--------+-------+-------+-------+-------+-------+

![](media/image107.png){width="2.8652777777777776in"
height="1.0708333333333333in"}

\(a\) (b)

[]{#__RefHeading___Toc506348237 .anchor}Figure 87 CCD images of
cross-sections of deionized water droplets deposited onto CdTe thin-film
after CdCl~2~ treatment (a) and subsequent etching in Br~2~:CH~3~OH (b).
In (a) the drop extends beyond the field of view to the right due to
greater surface wetting.

## **5.4 Conclusion**

Contact wetting is a simple and effective tool that is able to detect
changes in wetting angles and surface energies resulting from post
deposition processing. Crystallization of ITO is detectable by increases
in polar energies. As-deposited CdTe seeks to satisfy surface bonds
through reaction with air ambient. Significant native oxidation of PVD
(111) CdTe occurs within hours upon exposure to the atmosphere and can
be observed by increases in wetting angles. Dry air storage of CdTe
prior to CdCl~2~ heat treatments forms CdO, which has been correlated
with improved device performance while films kept in saturated air
produce CdTe~2~O~5~. The formation of CdTe~2~O~5~ may hinder p-type
device behavior due to a limited ability to form Cd vacancies due to
disproportionate removal of Te from the lattice. Polar energies increase
as a result of oxides formed during CdCl~2~ treatment and decrease upon
removal via chemical etching.

# 6 References

[^1]: J.E. Phillips, R.W. Birkmire, B.E. McCandless, P.V. Meyers, W.N
    Shafarman, *phys. stat. sol. (b)* **31,** 31 (1996).

[^2]: B.E. McCandless, J.E. Phillips, J. Titus, Proc. 2^nd^ WCPEC, 448
    (1998).

[^3]: S.S. Hegedus, W.N. Shafarman, *Progress in Photovoltaics*,
    **12**(2-3), 155 (2004).

[^4]: B. E. McCandless, J. R. Sites, *Handbook of PV Science and
    Engineering,* J. Wiley and Sons, 617 (2003).

[^5]: A. Gilmore, V. Kaydanov, T. Ohno, D. Rose, S. Feldman, P. Erselv,
    Proc. 29^th^ IEEE PVSC, 604 (2002).

[^6]: 1998 Annual Report to NREL Subcontract No. XAK-7-17609-01, 69-73.

[^7]: V.G. Karpov. A.D. Compaan, D. Shvydka, *Phys. Rev. B* **69,** 435
    (2004).

[^8]: Y. Marfaing, *Thin Solid Films*, **387,** 123 (2001).

[^9]: S. Hegedus, *Progress in Photovoltaics* **5(5),** 305 (1997).

[^10]: K. Mitchell, A. Fahrenbruch, R. Bube, *J. Appl. Phys.*
    **48**(10), 4365 (1977).

[^11]: R. Sudharsanan, A. Rohatgi, *Solar Cells* **31,** 143 (1991).

[^12]: D. Fardig, J. Phillips, Proc. 22^nd^ IEEE PVSC, 1146 (1991).

[^13]: D. Goren, G. Asa, Y. Nemirovsky, *Solar Energy Materials & Solar
    Cells*, **60**(1), 367 (2000).

[^14]: M. Gloeckler, J. Sites, *J. Applied Physics* **95**(8), 4438
    (2004).

[^15]: S. Hegedus, D. Ryan, K. Dobson, B. McCandless, D. Desai, *Proc.
    Mat. Res. Soc. Symp.* **763,** 447 (2003).

[^16]: S. Hegedus, W. Shafarman, *Progress in Photovoltaics*
    **12(2-3),** 155 (2004).

[^17]: P. Meyers, J. Phillips, Proc. 25^th^ IEEE PVSC*,* 789 (1996).

[^18]: S. Hegedus, B. McCandless, R. Birkmire, Proc. 28^th^ IEEE PVSC*,*
    535 (2000).

[^19]: S. Hegedus, B. McCandless, *Solar Energy Materials and Solar
    Cells,* in press 2005.

[^20]: X. Liu, J. Sites, *J. Applied Physics* **75**(1), 577 (1994).

[^21]: R. Crandall, *J. Applied Physics* **54**(12), 7176 (1983).

[^22]: Except we used no midgap states and a CdTe hole density of 1E14
    cm^-3^.

[^23]: A. Mondal, B. McCandless, R. Birkmire, *Solar Energy Mat & Solar
    Cells* **26**(3), 181 (1992).

[^24]: B. E. McCandless, R.W. Birkmire, W.A. Buchanan, S. Fields, G.M.
    Hanket, Proc. 29^th^ IEEE PVSC, 547 (2002).

[^25]: P. Meyers, C. Liu, T. Frey, US Patent 4,710,589 (1987).

[^26]: B.E. McCandless, H. Hichri, G.M. Hanket, R.W. Birkmire, Proc.
    25^th^ IEEE PVSC, 781 (1996).

[^27]: A. Fischer, C. Narayanswamy, D.S. Grecu, E. Bykov, S.A. Nance,
    U.N. Jayamaha, G. Contreras-Puente,

    A.D. Compaan, M.A. Stan, A. Mason, Proc. 25^th^ IEEE PVSC, 921
    (1996).

[^28]: Y. Qu, P.V. Meyers, B.E. McCandless, Proc. 25^th^ IEEE PVSC, 1013
    (1996).

[^29]: B.E. McCandless, R.W. Birkmire, Proc. 16^th^ European PVSEC, 349
    (2000).

[^30]: R.B. Bird, W.E. Stewart, E.N. Lightfoot, *Transport Phenomena*,
    John Wiley & Sons, New York (1996).

[^31]: O. Knacke, O. Kubaschewski, K. Hesselmann, *Thermochemical
    Properties of Inorganic Substances,* Springer-Verlag (1997).

[^32]: C.O. Bennett, J.E. Myers, *Momentum, Heat and Mass Transfer,*
    McGraw-Hill Book Company, New York, 590 (1982).

[^33]: A. E. J. De Nooy, M. Pagliaro, H. Van Bekkum, A. C. Besemer,
    *Carbohydr. Res.*, **304**, 117 (1997).

[^34]: V. Kumar, T. Yang, *Carbohydrate Polymers*, **48**, 403 (2002).

[^35]: N. Stratieva, R. Stefanov, G. Pchelarov, M. Stoev, A. Katerski,
    A. N. Tiwari, M. Igalson, Proc. of 2^nd^ World Conf. PVSEC, 1035
    (1998).

[^36]: K.D. Dobson, P.D. Paulson, B.E. McCandless and R.W. Birkmire, MRS
    Spring Meeting, B3.1 (2003).

[^37]: B. E. McCandless and K. D. Dobson, *Solar Energy*, in press 2005.

[^38]: S. Mu, J. Kan, *Synth. Met.*, **92,** 149 (1998).

[^39]: M. R. Hoffman, S. T. Martin, W. Choi, D. W. Bahneman, *Chem.
    Rev.*, **95**, 69 (1995).

[^40]: P.D. Paulson, M.W. Haimbodi, S. Marsillac, R.W. Birkmire, W.N.
    Shafarman, *J. Appl. Phys.* **91**(12), 10153 (2002).

[^41]: S. Marsillac, P.D. Paulson, M.W. Haimbodi, R.W. Birkmire, W.N.
    Shafarman, *Appl. Phys. Lett.* **81**(7), 1350 (2002).

[^42]: W.N. Shafarman, S. Marsillac, T. Minemoto, P.D. Paulson, R.W.
    Birkmire, Proc. 3^rd^ World Conf. PVSEC, 519 (2003).

[^43]: G.M. Hanket, U.P. Singh, E.Eser, W.N. Shafarman, R.W. Birkmire,
    Proc. 29^th^ IEEE PVSC, 567 (2002).

[^44]: S. Nishiwaki, N. Kohara, T.Negami, T. Wada, *Jpn. J. Appl. Phys.*
    **37**(2), L71 (1998).

[^45]: N. Kohara, S. Nishiwaki, Y. Hashimoto, T. Negami and T. Wada,
    *Sol. Energy. Mater & Sol. Cells.* **67,** 209 (2001).

[^46]: J. Keranena, J. Lua, J. Barnarda, J. Sterner b, J. Kessler, L.
    Stolt, Th.W. Matthes, E. Olsson, *Thin Solid Films* **387**(1-2), 80
    (2001).

[^47]: T.Nakada, H.Ohno, T.Watanabe, H. Nakazawa, M.Matsui and A.Kunika,
    *Sol. Energy Mater. & Sol. Cells* **49**, 285(1997).

[^48]: P.D. Paulson, R.W. Birkmire, W.N. Shafarman, *J. Appl. Phys.*
    **94**, 879 (2003).

[^49]: T. Gödecke, T. Haalboom and F. Ernst, *Z. Metallkd*. *Journal*
    **91**, 8 (2000).

[^50]: M. Turcu and U. Rau, *J. Phys. Chem. Solids*, **64**, 1591
    (2003).

[^51]: B. Sang, W.N. Shafarman, and R.W. Birkmire, Proc. 29^th^ IEEE
    PVSC, 632 (2002).

[^52]: 2003 Annual Report to NREL contract no. ADJ-1-30630-12.

[^53]: B. E. McCandless and W. N. Shafarman, Proc. 3^rd^ World Conf.
    PVSEC, (2003).

[^54]: O. Nast and A.J. Hartmann, J. Appl. Phys. **88**, 716 (2000).

[^55]: B.E. McCandless and JK.D. Dobson, *Sol. Energy* **77**(6), 839
    (2004).

[^56]: B.E. McCandless, and J.E. Sites, *Handbook of Photovoltaic
    Science and Engineering,* edited by S.S. Hegedus and A. Luque
    (Wiley, England, 2003).

[^57]: B.E. McCandless, and R.W. Birkmire, *Solar Cells* **31**, 527
    (1991).

[^58]: B.E. McCandless, L.V. Moulton and R.W. Birkmire, *Progress in
    Photovoltaics* **5**(4), 249 (1997).

[^59]: B.E. McCandless, Mat. Res. Sypm. Proc. **668**, H1.6.1 (2001).

[^60]: B.E. McCandless and R.W. Birkmire, Proc. 31^st^ IEEE PVSC,
    (2005).

[^61]: B.E. McCandless and S.S. Hegedus, (unpublished).

[^62]: S. Wei and S.B. Zhang, *Phys. Rev. B* **66**, 155 (2002).

[^63]: B.E. McCandless, I. Youm, and R.W. Birkmire, *Progress in
    Photovoltaics* **7**(1), 21 (1999).

[^64]: B.E. McCandless, S.S. Hegedus, R.W. Birkmire, D.Cunningham, *Thin
    Solid Films* **431-432**, 249 (2003).

[^65]: 2000 Annual Report to NREL, contract no. DE-AC 36-98-GO10337, 68.

[^66]: K. Durose, S.E. Asher, W. Jaegermann, D. Levi, B.E. McCandless,
    W. Metzger, H. Moutinho, P.D. Paulson, C.L. Perkins, J.R. Sites, G.
    Teeter and M. Terheggen, *Progress in Photovoltaics* **12**(2-3),
    177 (2004).

[^67]: J.S. Kim, R.H. Friend, F. Cacialli, *Synthetic Metals*
    **111-112**, 369 (2000).

[^68]: T. Young, Phil. *Trans. R. Soc. London* **9**, 255, 1805.

[^69]: M.Chasse and G.G. Ross, Nucl. Instru. Methods in *Phys. Res. B*
    **193**, 842 (2002).

[^70]: C.D. Lockhande, A. Barkschat and H. Tributsch, *Solar Energy
    Mat.* & *Solar Cells* **79**, 295 (2003).

[^71]: S Wu, *Polymer Interface and Adhesion* (Marcel Dekker, New York,
    178 (1982).

[^72]: C.W. Extrand and Y. Kumagai, *J. Colloid Interface Sci.* **191**,
    379 (1997).

[^73]: B. Giles, *Thin Film Characterization by Advanced X-Ray
    Diffraction Techniques,* edited by G. Cappuccio, F. Wang, A.
    Schwartzman, A.L. Fahrenbruch, R.R. Sinclair, M.L. Terranova
    (Laboratori Nazionali Di Fracasti: Frascati, Italy, 1996, pp.
    177-204.

[^74]: F. Wang, A. Schwartzman, A.L Fahrenbruch, R. Sinclair, R. H.
    Bube, *J. Appl. Phys*. **62**, (1987).