High Efficiency and High Rate Deposited Amorphous Silicon-Based Solar
Cells

PHASE II Third Quarter

Technical Progress Report

March 1, 2003 to May 31, 2003

**NREL Subcontract No. NDJ-2-30630-08**

Subcontractor: The University of Toledo

Principal Investigator: Xunming Deng

> Department of Physics and Astronomy
>
> University of Toledo, Toledo, OH 43606
>
> \(419\) 530-4782
>
> dengx@physics.utoledo.edu

Contract technical monitor: Dr. Bolko von Roedern

**Amorphous Silicon Deposited Using Trisilane as Precursor**

Contributors: W. Du, X. Yang and X. Deng

# Abstract

We explored the deposition of hydrogenated amorphous silicon (a-Si:H)
and hydrogenated nanocrystalline silicon (nc-Si:H) using trisilane
(Si~3~H~8~) gas as the silicon precursor in a plasma enhanced CVD
process, to improve our understanding of the growth chemistry and the
dependence of Si:H properties on the use of precursors. The film study
shows that the transition from amorphous to amorphous/nanocrystalline
mixed phase with increasing hydrogen dilution R
((\[H~2~\]/\[Si~3~H~8~\]) or (\[H~2~\]/\[Si~2~H~6~\])) is more gradual
for Si~3~H~8~-deposited Si:H as compared with disilane (Si~2~H~6~)
deposited Si:H and for the former the transition occurs at a higher
hydrogen dilution level R. The deposition rate of Si:H with different
hydrogen dilution is more dependent on the atomic ratio of H to Si
rather than the ratio of molecular hydrogen and silicon-containing
precursors. Open circuit voltages (V~oc~) of a-Si:H single-junction
n-i-p solar cells versus hydrogen dilution is used as a measure of
nanocrystalline formtion in the i-layer. Highest V~oc~ is obtained when
a-Si:H films are deposited using Si~3~H~8~ at R\~200 while highest V~oc~
is obtained for Si~2~H~6~ at R\~100.

# Introduction

Various silicon precursors, including SiH~4~, Si~2~H~6~, SiF~4~, and
SiH~2~Cl~2~, have been used for the deposition of a-Si:H (or a-Si) in a
plasma enhanced chemical vapor deposition (PECVD, rf 13.56 MHz) process.
Among these different precursors, SiH~4~ and Si~2~H~6~ were found to
produce a-Si:H materials with the highest photovoltaic quality. Compared
with SiH~4~, Si~2~H~6~ has higher dissociation rate in plasma and
results in different concentrations of reactive species such as SiH~3~
in the plasma. Such a different concentrations of reactive SiH~3~
species, believed to be responsible for the growth of high quality
a-Si:H, leads to different deposition rate and different optimal level
of hydrogen dilution for the growth of high quality a-Si:H.

When these silicon precursors are sufficiently diluted in hydrogen
during PECVD process, nc-Si:H (or nc-Si), having grain size in the order
of magnitude of 10nm, is often obtained. In the intermediate level of
hydrogen dilution, a mixed phase with coexisting a-Si:H and nc-Si:H is
usually obtained where nanometer sized Si crystals remain embedded in
amorphous Si matrix. It is generally found that a-Si:H with highest
stability under extended light illumination accompanied with reasonably
high photosensitivity is often obtained at a hydrogen dilution level
right below the transition from amorphous to mixed phase. A-Si:H
deposited under these conditions are often referred to as
"protocrystalline Si", or "edge Si" materials. The range of optimal
hydrogen dilution for the deposition of high-quality protocrystalline Si
depends on the selection of Si precursor and other deposition conditions
and is generally quite narrow. It also changes with the thickness of the
a-Si:H film as well as the types of substrates. The sharp nature of
transition region hinders reproducibility of the materials deposited
within. It is desirable to have a silicon precursor that offers a wider
range of hydrogen dilution for the deposition of high quality
protocrystalline Si material.

Trisilane, Si~3~H~8~, is another member of the silicon-hydrogen compound
gasses. It has a different dissociation rate than SiH~4~ and Si~2~H~6~
in an rf plasma. It is expected that Si~3~H~8~ could lead to a higher
concentration of SiH~3~ species in plasma when the gas is diluted in
appropriate amount of hydrogen. This may lead to higher quality a-Si:H
and/or wider window for the deposition of high quality a-Si:H. A
systematic study of Si~3~H~8~ deposition of Si:H could lead to a deeper
understanding of the nature of transition from amorphous to mixed and
finally to nanocrystalline phase of Si:H during PECVD growth under high
hydrogen dilution. Such a study may also lead to a-Si:H material with
improved stability under extended illumination. In this work, the growth
of Si:H near the transition region using Si~3~H~8~ and Si~2~H~6~ as the
precursor gasses is studied. This comparative study is carried out by
depositing and characterizing both single layer undoped films and solar
cell devices.

# 

# Experiment details

A-Si:H films and devices are grown in a multi-chamber PECVD system with
a load lock chamber. In such a study, it is important to make sure that
the measurement of the Si~3~H~8~ and Si~2~H~6~ gas flows is reliable.
For this reason, the Si~3~H~8~ and Si~2~H~6~ gasses, both obtained from
Voltaix, Inc., are introduced into the PECVD chamber via a common mass
flow controller. The gas factors of the mass flow controller for
Si~3~H~8~ and Si~2~H~6~ gasses were calibrated and verified by
pressurizing the deposition chamber using Si~3~H~8~ and Si~2~H~6~ gasses
separately, both through the same mass flow controller. In addition, the
deposition of Si:H films and devices using Si~3~H~8~ and Si~2~H~6~ were
carried out alternatively to avoid any unwanted effect due to the
systematic and unknown drift, if any, of the deposition equipment and
device processing procedures.

While hydrogen dilution is changed during the growth of these films, the
other deposition conditions are kept to be the same as the standard
i-layer deposition conditions used for the deposition of high-quality
a-Si:H solar cells in the same PECVD system. These conditions: substrate
temperature T~s~ at 150ºC, chamber pressure at 0.6Torr and rf power at
25mW/cm^2^, are kept unchanged for films deposited using both Si~3~H~8~
and Si~2~H~6~. The hydrogen dilution R was varied from 71 to 286 and
from 75 to 149 using Si~3~H~8~ and Si~2~H~6~ respectively. Various
substrates, including quartz, glass (Corning 1737), c-Si wafer, and
stainless steel foil, are used for Si:H deposition for various optical,
electrical and structural characterization. The film thicknesses range
from 250nm to 500nm.

A-Si:H solar cells with n-i-p structure were deposited on Al/ZnO coated
stainless steel. Intrinsic-layer deposition condition was the same as
described above. The window p-layer was deposited at low T~s~, about
70-100 ºC, and high hydrogen dilution. ITO was sputtered on the p-layer
as transparent conductive coating. Light soaking was done under white
light at 1 sun intensity and the ambient temperature was \~ 40-50C.

Infrared absorption spectroscopy was measured for Si:H films deposited
on c-Si substrate in the ranges from 400cm^-1^ to 800 cm^-1^ and from
1800cm^-1^ to 2400 cm^-1^ and with a resolution of 0.33 cm^-1^. The
integrated strength for the wagging mode at 640cm^-1^ was used to obtain
the H content (at.%). The stretching modes at 2000cm^-1^ and 2100 cm^-1^
were used to obtain the microstructure factor R\* in the film. Gaussian
distribution was used to fit the absorption peaks. Transmission spectra
were recorded using a Cary 5 spectrophotometer. Scanning range was from
400nm to 2200nm. Thickness was calculated using the method described in
the literature.^\[1\]^ Raman scattering was performed to characterize
the structure of Si:H films. 488nm laser line from a water-cooled Ar^+^
laser was used for the excitation. Photoconductivity was measured when
the samples were annealed in a test chamber purged with N~2~. A pair of
Ag electrodes, \~2 mm apart, were painted on the films. Samples are
heated to 180ºC and the temperature dependence of dark conductivity is
measured during the cooling process, from 180ºC to 50ºC. Activation
energy was measured from a semilog plot (log *vs* 1000/T).
Photoconductivity was measured using a ELH lamp at 100 mW/cm^2^
intensity at room temperature (25ºC).

# 

# RESULTS AND Discussions 

Figure 1 shows the deposition rate of intrinsic layers as a function of
hydrogen dilution R. Film thickness was calculated from the transmission
spectrum. The deposition rate of Si~3~H~8~ samples is higher than that
of Si~2~H~6~ samples when the two films are deposited with the same R.
As R is increased from 75 to 285 for Si~3~H~8~ samples, deposition rate
is decreased from 0.45 Å/sec to 0.2 Å/sec. Lower deposition rate is
observed for Si~2~H~6~ samples with the same R. In each series, smaller
deposition rate is observed at higher dilution R. This decrease in
deposition rate often accompanies with formation of nanocrystalline
phase which is a more ordered structure. The increase in hydrogen
dilution provides the required atomic hydrogen to etch out the
disordered portion of the growing a-Si:H films and also provide the
hydrogen coverage on the growing surface. Thus the deposition rate
decreases and simultaneously these factors induce the propagation of
nanocrystalline network in the films. It is additionally observed that,
R~t~, the hydrogen dilution level at which the transition from amorphous
to an amorphous/nanocrystalline silicon mixed phase occurs, is higher
for a-Si:H films deposited using Si~3~H~8~. R~t~ is around 200 for
Si~3~H~8~ while R~t~ is around 150 for Si~2~H~6~. The deposition rate is
also plotted against the atomic dilution ratio, R~a~=\[H\]/\[Si\], which
is the ratio of number of hydrogen and silicon atoms in the gas. Since
R~a~ is identical to R for Si~2~H~6~, only deposition rate dependence on
R~a~ for Si~3~H~8~ samples is added in figure 1, as solid triangle
symbols. It is further observed that the deposition rate *vs* R~a~
curves are approximately the same for both gasses. Therefore, the
deposition rate is more influenced by the ratio of the total number of
Si atoms to hydrogen atoms directed into the chamber rather than the
number of Si-containing molecules.

![](media/image1.wmf)

Figure 1. Deposition rates of intrinsic a-Si:H film deposited under the
conditions of RF power =25mW/cm^2^ and T~s~=150ºC. Square symbols
represent Si~2~H~6~ series and diamond symbols represent Si~3~H~8~
series; The solid triangles represent the relationship between
deposition rates and R~a~=\[H\]/\[Si\] for Si~3~H~8~ series.

It is widely known that high-quality a-Si:H material is often deposited
near the transition region from amorphous to nanocrystalline via a
hydrogen dilution mediated path. A material deposited near such a
transition region is sometimes referred to as protocrystalline or "edge"
material. In order to find R~t~ for both Si~2~H~6~ and Si~3~H~8~, a-Si:H
films are deposited on a variety of substrates at different R.
Si~3~H~8~-deposited films in this study are deposited with R (H~2~
(sccm):Si~3~H~8~(sccm)) of: 71 (50:0.7); 107 (75:0.7); 143 (100:0.7);
179 (100:0.56), 213 (100:0.47) and 286 (100:0.35). Si~2~H~6~-deposited
films in this study are deposited with R (H~2~(sccm):Si~2~H~6~(sccm)) of
75 (75:1); 100 (100:1); 125 (100:0.8); and 149 (100:0.67). Figures 2A
and 2B show the Raman spectra of two series of films deposited using
Si~3~H~8~ and Si~2~H~6~, respectively. Table 1 shows the volume fraction
of nc-Si:H phase calculated using approaches described in the
literature^\[2\]^. Figure 2A shows that at R=143, the material is still
amorphous while at R=179 formation of a-Si:H/nc-Si:H mixed phase starts.
The volume fraction of the nc-Si:H phase calculated from the peak
intensity is 2.8%. Therefore, R~t~ is between 143 and 179 for Si~3~H~8~.
For Si~2~H~6~-deposited samples, Figure 2B indicates that R~t~ is
between 100 and 125, lower than that of Si~3~H~8~. It should be noted
that Sample GD1006 with R=179 is thicker than other films such as
GD1002. Under a high R, thicker a-Si:H is more likely to have nc-Si:H
phase ^\[3\]^ It is likely that R~t~ for Si~3~H~8~-deposited films is
higher than 179 if the film thickness would be around 170nm (as in
GD1002), which is typical for the i-layer of a-Si:H top component cell
in a triple-junction solar cell. As to be described later, solar cells
made with Si~3~H~8~ at R=179 showed high V~oc~, indicating that R~t~ is
greater than 179 for Si~3~H~8~-deposited films with around 170nm
thickness. These Raman studies suggest that R\~100 and R\~180, for
Si~2~H~6~ and Si~3~H~8~, respectively, may be optimum for deposition of
a-Si:H solar cell intrinsic layers with thickness around150 to 200 nm.

The films described in Figure 2 and Table 1 were also evaluated using
Raman with excitation light through the quartz substrates. Such a
measurement reveals the structure of the film grown during initial stage
of the deposition. Raman scattering measurement show that the a-Si:H
materials deposited on glass or quartz substrates during the initial
stage of deposition is amorphous for all these samples described in
Table 1 and Fig. 2.

![](media/image2.wmf)![](media/image3.wmf)

\(a\) (b)

Figure 2 Raman spectra of the series of sample deposited using Si~3~H~8~
(2a) and Si~2~H~6~ (2b). In Fig. 2a, curves from bottom to top are for
Si:H deposited with R=143, 179, 213, and 286. In Fig. 2b, curves from
bottom to top are for a-Si:H deposited with R=75; 100; 125; and 150.

Table 1. Volume fraction of nanocrystalline silicon phase estimated from
Raman spectra for two series of sample. Thicknesses of samples are also
included in the Table.

  ----------- ----------- ----------- --------- -------------- --------------
  Sample No   Gas         Gas flow    Ratio     Thickness      nc --Si:H
              precursor                                        volume frac.

                          (sccm)      R         (nm)           (%)

  GD1003      Si~2~H~6~   1:75        75        176            N.A.

  GD1002      Si~2~H~6~   1:100       100       165            N.A.

  GD1004      Si~2~H~6~   0.8:100     125       195            13.6

  GD1005      Si~2~H~6~   0.67:100    149       209            37.8

  GD1010      Si~3~H~8~   0.7:100     143       234            N.A.

  GD1006      Si~3~H~8~   0.56:100    179       340            2.8

  GD1009      Si~3~H~8~   0.47:100    213       261            17.7

  GD1007      Si~3~H~8~   0.35:100    286       273            57.2
  ----------- ----------- ----------- --------- -------------- --------------

Activation energy E~a~ was calculated from temperature dependence of
dark conductivity using =~0~exp(-E~a~/kT). For undoped or n-type a-Si:H,
E~a~ is approximately the energy difference between the Fermi level and
the conduction band edge. Figure 3 shows the activation energy as a
function of hydrogen dilution R for both Si~3~H~8~- and
Si~2~H~6~-deposited samples. For Si~3~H~8~-deposited samples, E~a~ is
decreased from 0.85 eV at low R to 0.5 eV at high R. This decrease is
believed to be due to the narrowing of the bandgaps of the deposited
materials when there is a higher volume fraction of nc-Si:H phase in the
sample. Similar trend is observed for Si~2~H~6~ deposited films. A major
difference between the two series of films is that E~a~ for Si~2~H~6~
samples decreases rapidly when R is increased to 150 while E~a~ for
Si~3~H~8~ samples decreases gradually, over a larger range of R, when R
is increased to \~250. The gradual change in E~a~ with increasing R may
suggest that the range of R for the deposition of a-Si:H/nc-Si:H mixed
phase is broader. The exact nature for such a difference between
Si~3~H~8~ and Si~2~H~6~ is unclear.

![](media/image4.wmf)

Figure 3 Activation energy of a-Si:H films made using Si~2~H~6~ and
Si~3~H~8~ as a function of hydrogen dilution R.

Fourier transform infrared spectroscopy (FTIR) was used to measure the
content of bonded hydrogen in these Si:H films as well as the amount of
hydrogen in different bonding environment. A microstructure factor, R\*,
defined as the fraction of hydrogen in dihydride form among all hydrogen
incorporated in the film, is calculated from the integrated strength of
the 2100cm^-1^ peak and 2000cm^-1^ peak (R\*= I (2100 cm^-1^)/ \[I (2100
cm^-1^)+I(2000 cm^-1^)\]).^\[4\]^ The microstructure factor R\* is
plotted against dilution R for Si~3~H~8~ and Si~2~H~6~ deposited films
in Fig. 4.

Figure 4 shows that for both series of films, R\* decreases initially
and increases again at higher R. The initial decrease is associated with
the reduction of undesirable micro-voids in the films. The increase in
R\* at higher dilution R is believed to be associated with the nc-Si:H
formation\--hydrogen at grain boundaries bonds to Si mostly in dihydride
form. In addition to these general trends, there are noticeable
difference in the R\* vs R between the two series of films. The increase
in R\* occurs at a higher R for Si~3~H~8~-deposited films, indicating
that R~t~ is higher for Si~3~H~8~. It is interesting to notice that for
films deposited without hydrogen dilution (R=0), R\* is lower for
Si~3~H~8~ films. This is unexpected because the atomic ratio of H to Si
is lower in Si~3~H~8~ than Si~2~H~6~. The lower R\* (at R=0) for
Si~3~H~8~ suggests that Si~3~H~8~ is dissociated in plasma into a form
that is more desirable for the formation of dense a-Si:H. However, one
needs to be cautious in drawing a general conclusion based on these
results since other deposition conditions such as rf power, chamber
pressure, substrate temperature may impact R\* at R=0.

![](media/image5.wmf)

Figure 4 Microstructure factor R\* as a function of hydrogen dilution R.

Bonded hydrogen content in the Si:H films is estimated from the
integrated peak strength of the Si-H wagging mode (640 cm^-1^). Figure 5
shows the H content in the a-Si:H films for both series. For Si~2~H~6~
samples, H content is approximately 15-20 at.% at R\~100. For Si~3~H~8~
samples, H content is approximately 20% at R\~150, and decreases down to
6% at R\>250. The absolute value of H content depends on the calibration
factor used and the values in Figure 5 may not be accurate in the
absolute scale. However, relative change in H content is believed to be
reproducible.

![](media/image6.wmf)

Figure 5. H content estimated from FTIR spectra as a function of H~2~
dilution R

Figure 6 shows photosensitivity of these a-Si:H films, which is the
ratio of photoconductivity under 1 sun illumination and the
dark-conductivity. The photosensitivity decreases with increased volume
fraction of nc-Si:H phase due to the defects near the grain boundaries.
The highest photosensitivity, 1×10^6^, occurs at R\~100, or R~t~, for
Si~2~H~6~ samples. The photosensitivity then decreases with increasing
R, indicating the increase of nanocrystalline fraction in the material.
The highest photosensitivity is nearly the same for the two series of
samples. However, the decrease in photosensitivity at R\>R~t~ is more
gradual for Si~3~H~8~, similar to the gradual decrease in its activation
energy.

![](media/image7.wmf)

Figure 6. Photosensitivity of a-Si:H films deposited using Si~2~H~6~ and
Si~3~H~8~ as a function of hydrogen dilution.

Results in Figures 3-6 indicate that when Si~3~H~8~ is used for the
deposition of Si, the transition from amorphous to nanocrystalline is
more gradual as compared with that for Si~2~H~6~. In other words, the H
dilution range for mixed phase is broader when Si~3~H~8~ is used.
Material with optimal optoelectronic performance is grown with R\~180 at
T~s~\~150 ^o^C when Si~3~H~8~ is used. These materials have similar
properties to those materials deposited using Si~2~H~6~ with R\~100.

In order to further verify the findings from this materials study, two
series of solar cells were fabricated using Si~3~H~8~ and Si~2~H~6~ with
different hydrogen dilution R for the i-layers. When all other device
parameters are kept the same, V~oc~ of the solar cells are directly
related to the i-layer bandgap. When appropriate level of R is used and
the Si:H is near the transition from amorphous to mixed phase, V~oc~ is
the highest. At higher R, V~oc~ drops due to the formation of nc-Si:H in
the i-layer. Figure 7 shows the V~oc~ of solar cells as a function of R
for i-layers made using Si~3~H~8~ and Si~2~H~6~ at different R. For
Si~2~H~6~ devices, V~oc~ reaches its highest value (1.014V) at R\~100,
while for Si~3~H~8~ devices, V~oc~ reaches its highest value (1.010V) at
R\~180. These results on solar cell devices are consistent with the
findings in the materials study of single-layer films deposited on
glass. Both series of solar cells, with \~200nm thick i-layers, were
illuminated with white light at 1 sun intensity. After 100 hours of
light soaking, Si~3~H~8~ devices degraded 11.7% and Si~2~H~6~ devices
degraded 13.2%. These degradations are considered to be the same within
experimental variations.

![](media/image8.wmf)

Figure 7 Open circuit voltage versus H~2~ dilution for Si~2~H~6~ and
Si~3~H~8~ devices.

# SUMMARY

Si:H films and solar cells deposited using Si~3~H~8~ and using Si~2~H~6~
are compared. At the same hydrogen dilution, Si~3~H~8~ leads to higher
deposition rate. However, at the same atomic ratio of hydrogen and
silicon, the deposition rate of Si:H using Si~3~H~8~ and Si~2~H~6~ is
approximately the same. The deposition rate is more influenced by the
ratio of the total number of Si atoms to hydrogen atoms introduced into
the chamber rather than the number of Si-containing molecules.
Measurement of activation energy, microstructure and photosensitivity
indicate that

1\) the transition from amorphous to amorphous/nanocrystalline mixed
phase occurs at a higher hydrogen dilution level when Si~3~H~8~ is used;
and

2\) the transition from amorphous to mixed phase is more gradual for
Si~3~H~8~ as compared with Si~2~H~6~.

When appropriate level of R is used and the a-Si:H is near the
transition from amorphous to mixed phase, V~oc~ is the highest. At
higher R, V~oc~ decreases due to the formation of nc-Si:H in the
i-layer. For Si~2~H~6~ devices, V~oc~ reaches its highest value (1.014V)
at R\~100, while for Si~3~H~8~ devices, V~oc~ reaches its highest value
(1.010V) at R\~180. These results on solar cell devices are consistent
with the findings in the materials study of single-layer films deposited
on glass.

# References

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