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the university of north carolina

[at]{.smallcaps}

[chapel hill]{.smallcaps}

Department of The University of North Carolina at Chapel Hill

Physics and Astronomy CB# 3255, Phillips Hall

College of Arts and Sciences Chapel Hill, N.C. 27599-3255

Tel: (919) 962-5002

Email: daxing@physics.unc.edu Fax: (919) 962-0480

Oct 31, 2003

Dr. Bolko von Roedern

NREL

1617 Cole Blvd.

Golden, CO 80401

Dear Bolko,

Here is the Phase II Annual Technical Report for the period of
Oct/1/02-Sept/31/03 on NREL Subcontract No. ADJ-1-30630-09 with the
University of North Carolina at Chapel Hill.

Sincerely,

Daxing Han

cc: Carolyn Lopez (NREL)

Donna Braxton (UNC)

Table of Content

I. Raman study of thin films from amorphous-to-nanocrystalline Si
prepared by HW-CVD.....2

II\. Micro-Raman Studies of Mixed-Phase Hydrogenated Silicon Solar
Cells......................11

III\. Toward understanding the degradation without light soaking in
hot-wire a-Si:H thin films and solar
cells............................................................................................16

IV\. Is there a correlation between the median component in Raman
spectrum and the (111) side peak in
XRD?..........................................................................................\...22

V. PL study of growth rate dependence on solar cells made by
BPSolar............................23

VI\. Review Article and publication list in
2003.........................................................25

Phase II Annual Technical report Oct/1/02-Sept/30/03

NREL Subcontract ADJ-1-30630-09 with the University of North Carolina at
Chapel Hill

Bolko von Roedern, NREL Technical Monitor;

Carolyn Lopez, NREL Subcontract Associate;

Daxing Han, principal investigator; Jessica M. Owens, J.R.
Weinberg-Wolf, graduate students; and Keda Wang, co-P.I.

Submitted: Oct 31, 2003

**I. Raman study of thin films from amorphous-to-nanocrystalline Si
prepared by HW-CVD**

**I.1 introduction**

Raman spectroscopy is a sensitive tool that provides valuable structural
information about a-Si-based materials. In a-Si, all phonon modes of the
transverse acoustic (TA), longitudinal acoustic (LA), longitudinal
optical (LO), and transverse optical (TO) modes are Raman active. Thus,
Raman spectra from a-Si at room temperature yield a reasonable spectral
comparison with phonon density of states that are modified substantially
by small changes in short-range order (SRO). Theoretical calculations
found that the full width of half maximum (FWHM), the peak frequency of
the TO phonon mode (~TO~), and the relative intensity of the TA mode,
TA/TO, are sensitive to bond-angle variation () in an a-Si network.
Beeman et al. found a linear relationship between the width of the TO
mode and the spread in the mean bond angle in an a-Si network as

= 15+6 (1)

More recent calculations using 1000-atom configurations showed a similar
relation of =18.4 + 6.6 and a shift of the TO mode frequency toward
higher frequency as decreases following

~TO~ = -2.5 +505.5 (2)

for HV-polarized light. The authors further found that the TA/TO
intensity ratio decreases linearly with decreasing . For device-quality
a-Si:H, the TO mode centered at 480 cm^-1^ with FWHM 70 cm^-1^.
According to equation (1), the spread in mean bond angle is ≤10. In
c-Si, on the other hand, the Raman spectrum consists of a single sharp
TO mode with FWHM of 4 cm^-1^ at a frequency of 520 cm^-1^. All other
modes are not Raman active because of the symmetry in the fcc lattice.
However, when the c-Si grain size is as small as a few nm, the momentum
conservation will be relaxed and Raman active modes will not be limited
to the center of the Brillouin zone. Thus, the frequency could shift
from 520 cm^-1^ to 512 cm^-1^ with decreasing grain size from 10 to 3
nm. Below 3 nm, a crystalline-to-amorphous transition occurs. In
addition to the a-Si and c-Si phonon modes, the Raman TO mode from nc-Si
films contains a third intermediate peak centered around 500 ± 10 cm^-1^
which could be attributed to the contribution of grain boundaries or
other structures. With hydrogen dilution, Tsu et al. found chainlike
objects in PE-CVD films using high-resolution transmission electronic
microscopy (TEM) which they claim is evidence for intermediate ordering
in H-diluted films. They found that the a-Si TO band centered at 490
cm^-1^ (37 3 cm^-1^ full width) instead of 480 cm^-1^.

To determine the c-Si volume fraction from the Raman TO mode, one must
take into account the differences between a-Si and c-Si with respect to
scattering cross sections and optical absorption coefficients. Bustarret
et al. found that the ratio of the cross section for the
amorphous-to-crystalline phase for grain size \> 3 nm at the excitation
wavelength 514.5 nm can be written as

y(L) = 0.1 + exp\[-(L/250)\]. (3)

It is expected that y(L) will be 2.4 times smaller at 632.8-nm
excitation. Based on the above argument, we deduced the crystalline
volume fraction, X~c~, from the Raman TO mode using

X~c~ =(I~c~ + I~gb~)/\[I~c~ + I~gb~ +y(L)I~a~\], (4)

where I~c~ + I~gb~ and I~a~ are integrated intensities of the c-Si,
intermediate, and a-Si peaks, respectively; y(L) decreases from 1 to 0.4
when the grain size L grows from 3 nm to 30 nm. We found that when using
514.5-nm laser front excitation, X~c~ can be deduced in such a way to be
consistent with the results from x-ray diffraction (XRD) for HW-CVD
transition films. In previous works, we reported Raman, XDR, FTIR,
photoluminescence, and optical absorption results for HW-CVD films with
either varied hydrogen-to-silane ratio, R, with fixed silane flow rate
or varied substrate temperature, Ts, at fixed R = 3. We observed a blue
shift of the a-Si:H PL peak frequency, which was due to narrowing of the
band tail states by hydrogen dilution and a low-energy PL band in the
nc-Si films. The low-energy PL band originates from the c-Si g.b.
regions, and thus, we have attributed the intermediate component of the
Raman TO mode to the g.b. regions. The previous Raman results showed a
threshold of R 2 at Ts = 240^○^-250^○^C and Ts 200^○^C at R = 3 for the
structural transition from a- to nc-Si:H. Both the threshold growth
conditions of crystallinity and the crystalline silicon volume fraction
deduced from the Raman TO modes are consistent with the results from
X-ray diffraction. In this work, we have extended our deposition
parameter range to include not only the H-dilution ratio R, but also the
effect of silane flow rate and total gas pressure. For example, at a
fixed R, the individual SiH~4~ and H~2~ flow rates can be changed so
that one condition has a higher silane flow rate than the other. The
films deposited at the same R were not the same and depended on the
actual gas flow rates. This systematic study of the correlation between
material structure and deposition conditions can help us understand the
growth mechanism and adjust the growth conditions of nc-Si:H films.

**I.2. sample and Experimental**

We studied 125 intrinsic Si films by HW-CVD on Corning 1737 25 25 mm^2^
glass. All of the films were grown at the same substrate temperature, Ts
= 240^o^C. There were four groups with various silane flow rates of 3,
8, 16, and 22 sccm, respectively. The hydrogen flow rate for each group
was varied from 0 to 216 sccm. The total gas pressure was controlled by
adjusting the throttling valve position (TVP) at 45%, 50%, 60%, 70%, and
80%. The higher percentage leads to lower pressure. For the films with a
low silane flow rate of 3 sccm, the thickness was controlled in the
range of 1500-2000 Å, and the growth rate was in the range of 1-3 Å/s;
for the films with silane flow rates of 8 and 16 sccm, the thickness was
controlled in the range of 5000-6000 Å, and the growth rate was in the
range of 5-13 Å/s; for the films with a silane flow rate of 22 sccm, the
thickness was controlled in the range of 0.9-1.1 µm, and the growth rate
was in the range of 8-21 Å/s. Two films were made as thin as the
penetration depth of the 514.5-nm laser, 60-70 nm, and these were used
for calibrating the Raman spectra from back excitation.

Raman spectra were measured using an XY triple spectrograph equipped
with a LN2-cooled CCD detector. The data were taken at room temperature
using the 514.5-nm line of an argon-ion laser. The spectral resolution
was approximately 0.5 cm^-1^ in a frequency range of 100 to 1000 cm^-1^.
The frequency of the Raman lines was calibrated using the TO mode of a
(111) c-Si wafer at 520 cm^-1^. The power of the laser was kept below
the level that would thermally induce crystallization. The penetration
depth of the 514.5-nm light is \~60 nm for a-Si:H and is larger for
nc-Si:H. Since the Raman spectra reveal the structure of the thin layer,
measurements were also done from back-excitation for most samples. To
depress the scattering light from the substrate, the light spot was
focused on the film using a microscope. The Raman TO mode for all films
was deconvoluted into two Gaussian functions and one narrow Lorenzian
function to represent the three components of a-Si, intermediate, and
c-Si peaks. The c-Si grain size ranged from 7 to 30 nm, which was
obtained from XRD; the crystalline volume fraction, X~c~, was deduced
using equations (3) and (4) with 0.4 \< y(L) \< 0.85.

#### II.3. Results and discussion

In general, the structure change is a function of the ratio of
hydrogen-to-silane flows, R. Fig. 1(a-p) shows the crystalline silicon
volume fraction, X~c~, as a function of R for four groups of films at
SiH~4~ flow rates of 3, 8, 16, and 22 sccm and with the total gas
pressure varied with the throttling valve position (TVP) of 45%, 50%,
60%, 70%, and 80%. The solid and dotted lines indicate the X~c~ obtained
from the top and bottom layer, respectively. The X~c~ is found from zero
to \>90% (including X~g.b.~ from zero to \~30%) in the top layers. In
agreement with the previous works, there is a threshold R for the
structural transition from a- to nc-Si:H. The common features of the
curves in Fig.1(a-p) are: (a) Narrow structure transition region (R\~1)
from X~c~= 0 to ≥50%. We will analyze this transition region in detail
shortly. (b) The threshold R of crystallinity increases with increasing
SiH~4~ flow rates. For instance, the threshold is R≤2 for low SiH~4~
flow rates of 3 and 8 sccm, as shown in Figs. 1(a-g), but no
crystallinity has yet occurred at R=2 for high SiH~4~ flow rates of 16
and 22 sccm as shown in Figs. 1(i-p). (c) The X~c~ changes with the
total pressure in the chamber in the following way: the films easily
become nc-Si at the higher pressure (lower TVP) and lower silane flow
rates of 3 and 8 sccm as seen by comparing Figs. 1(a) to 1(d) and 1(e)
to 1(h); whereas, at high silane flow rates of 16 and 22 sccm, a lower
pressure (higher TVP) yields a higher X~c~ in the film as seen by
comparing Figs. 1(i) to 1(j) or 1(m) to 1(n). These surprising results
suggest that there is a "growth zone" for the crystallite, and it is our
intention to study many deposition parameters, and not just R, to
examine this further. We used back excitation to study the early growth
of the film. At the early stage of the growth, X~c~ is much lower than
in the late growth of the films as shown by the dotted lines in Fig 1.
Whereas, in Fig. 1(d) for the film with R = 72 and SiH~4~ = 3 sccm, we
observe roughly the same X~c~ value. This is because the film thickness
of 70 nm is close to the penetration depth of the 514.5-nm laser. For
all of the films in this study, the FWMHs of the a-Si network from the
bottom layers are in the range of 75-83 cm^-1^. These values are always
broader than the 67-80 cm^-1^ for the top layers. This implies that the
a-Si:H network near the glass substrate is less ordered compared to the
top surface layer.

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Fig. 1. X~c~ as a function of R at SiH~4~ flow rates of 3, 8, 16 and 22
sccm. The solid and dotted lines indicate the X~c~ was obtained from the
top or bottom layer of 60 nm, respectively.

![](media/image18.wmf){width="3.5972222222222223in"
height="3.5972222222222223in"}

![](media/image20.wmf){width="3.5972222222222223in"
height="3.5972222222222223in"}Fig. 2. Raman shift and their fittings
from the films with SiH~4~ flow of 3 sccm, TVP=45%, (a) R=1, and (b)
R=2. And the enlarged TO mode of the data in 2(c) and 2(d).

Figs. 2(a) and 2(b) show the Raman spectra and their fitting functions
for the films with SiH~4~ flow of 3 sccm, TVP=45%, and R=1 and 2,
respectively. The fittings for the TA mode peaked at 145 cm^-1^ and the
LA mode peaked at 330 cm^-1^, which are also shown. The signal between
the TO and LA modes is attributed to the LO mode. Comparing Fig. 2(a)
and 2(b), one can see that the intensity ratio of TA/TO decreased as the
film was more crystallized. This agrees qualitatively with the
theoretical calculations. More importantly, we obtain the structure
information from the TO mode according to equations (1) and (2). Figs.
2(c) and 2(d) show the enlarged TO mode between 450 to 550 cm^-1^ for
the same samples in Figs. 2(a) and 2(b). The band width of the TO mode,
FWHM, and their peak frequency, ~TO~, as a function of H-dilution ratio
are plotted in Figs. 3(a) and 3(b) for all the films in this study.
Generally, the FWHM decreases and the ~TO~ increases in the narrow
structural transition region and then reaches a saturation value. Notice
that the films with a high SiH~4~ flow rate of 22 sccm (the big crosses)
show the narrowest FWHM and lowest ~TO~. It is perhaps related to
column-like structures. The changes can be seen more clearly for the low
growth rate films with a SiH~4~ flow rate of 3 sccm which are
represented by solid dots in both Fig. 3(a) and 3(b).

![](media/image22.wmf){width="3.375in" height="3.375in"}
![](media/image23.wmf){width="3.375in" height="3.375in"}

Figs. 3.(a) FWHM and (b) ~a-Si~ of the TO mode as function of R. The
solid dots, dashes, open circles, and crosses represent the films with a
SiH~4~ flow rate of 3, 8, 16, and 22 sccm, respectively. There are two
regimes separated at R3: the fast changing regime when R \<3 and the
saturated regime when R\>3.

The improvement of SRO of the a-Si:H network in the a-to-nc-Si
transition region is clearly shown in both the changes of FWHM and ~TO~.
In Fig. 3(a), the FWHM decreases from 80 to 67 cm^-1^ when 0 ≤ R ≤ 3,
and then remains at 67 cm^-1^ when R \> 3. This indicates that the
spread in the mean bond angle decreased from 10.8 to 8.6 according to
equation (1). The corresponding frequency shift ~TO~ for the same films
in Fig. 3(b) is from 475 cm^-1^ to 480 cm^-1^, which indicates that the
decreased from 12.2 to 10.2 according to Eq. (2). The absolute values of
are relatively large from the calculation, but the 2 improvement of the
bond angle spread is consistent with the results from the decrease of
the band width. The improvement of SRO of the a-Si network by increasing
H-dilution is consistent with the observations of a narrowing of the
valence band tail in our previous work. For most of the
high-SiH~4~-flow-rate films, indicated by the open circles and big
crosses, the FWHM = 67 cm^-1^ does not change with R, but the ~TO~
spreads from 475 to 480 cm^-1^ at 0\<R\<3. Whereas, the FWHM is as
narrow as 58.5 cm^-1^ and the peak frequency as low as 472 cm^-1^ for
the films deposited with the highest growth rate of \~20 Å/s
(represented by the big crosses). The frequencies of the c-Si and g.b.
peaks for those films are also low (see Figs. 4 and 5). Since a
column-type growth occurs in such conditions, it implies that both the
~TO~ and the FWHM are sensitive to the column structures. To our
knowledge, there are no calculations of how the phonon density of states
changes with such structures. For the c-Si component, we generally found
that the width of the c-Si TO mode for the nc-Si:H films is broader than
that for the c-Si wafer, i.e., 8-10 cm-1 compared to 4 cm-1. There were
two exceptions for films with SiH~4~=3 sccm with TPV=70% and R=1.2 and
3. The reason is unclear why the FWHM was as narrow as 4 cm-1 for those
two films. However, the peak frequency of the c-Si TO mode, ~c-Si~,
changes when R increases. For the films with SiH~4~ flow rate of 3 and 8
sccm, we found a blue shift from 516 to 520 cm-1 in the transition
region, as shown in Fig. 4. This indicates that the crystallite size has
increased from \~7 nm to ≥ 10 nm. For the films with a high SiH~4~ flow
rate of 22 sccm, the big crosses indicate that the ~c-Si~ blue shift
occurs from 513 to 518 cm-1. Again, the relatively low frequency of the
c-Si TO mode could be due to the column-like structures.

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Fig. 4. ~c-Si~ as a function of R. Fig. 5. ~g.b~. as a function of R.

Finally, for all the nc-Si:H films, the Raman TO mode always contains an
intermediate component in addition to the a-Si and c-Si components. The
intermediate component was fitted well using a Gaussian function with
FWHM of 33 cm-1 for all the nc-Si:H films. The peak frequency of the
intermediate component, ~g.b~, as a function of R is plotted in Fig. 5.
One finds that the value spreads from 500 to 510 cm-1. For the
low-growth-rate films, there is a blue shift of the ~g.b~ from 500 to
507 cm-1 in the transition region and thenspreads between 505 to 507
cm-1 when R \>3. Regardless of whether the intermediate component
originates from the grain boundaries, small size (\<3 nm) crystallites,
or median-range ordered objects, the blue shift of the ~g.b~ can be
attributed to the growth of the c-Si grain size. One may argue that
Raman modes reflect SRO, and the phonon density of states is sensitive
to the local bond angles. However, the scattering intensity is related
to a coupling-parameter-weighted phonon density of states. Small
differences between peak positions could be attributed to changes in
median- or long-range interactions that are consequences of the coupling
parameter. We could not clarify the origin of this intermediate
component in this work. Whatever the origin is, one commonly takes the
third component into account as part of the c-Si volume fraction as seen
in equation (4). As shown above, the peak frequencies of all the three
components of the TO modes show blue shift with an increase of
H-dilution in the structural transition region. The compressive stress
could be higher in films with higher hydrogen dilution ratios (the film
easily peels off from the substrate), and the blue shift of the phonon
bands could be partially related to an increase of the film stress.

**I.4. Summary**

Raman scattering studies were carried out on the transition films from
amorphous-to-nanocrystalline silicon prepared by HW-CVD with varied
silane and hydrogen flow rates and the T-valve position. The
measurements were done for both front and back excitation. The TO mode
was deconvoluted into a broad Gaussian function for the a-Si:H and a
narrow Lorenzian function for the c-Si; a Gaussian function peaked at
between 500 to 510 cm-1 with FWHM of 33 cm-1 for the intermediate
component. We believe that the origin of the intermediate component is
the grain boundaries and the small-size (\< 3 nm) crystallites, and
attribute it into the total c-Si volume fraction. The c-Si volume
fraction depends not only on the H-dilution ratio, R= H~2~/SiH~4~, and
the film thickness, but also on the SiH~4~ flow rate and total gas
pressure which can be explained using the silane-depletion model. The
detailed calculation of gas depletion will be reported in a future work.

The most interesting finding is a narrow structural transition region of
R\~1, in which the short-range order of the a-Si:H network is improved:
the width of the TO mode decreased from 80 to 67 cm^-1^ and the peak
frequency, ~TO~, blue shift from 475 to 480 cm^-1^. These indicate that
the bond-angle variation, ,decreases from \~10˚ to \~8˚ when the
crystallites start to grow. Furthermore, the crystallite grain size
increases from \~7 nm to ≥ 10 nm in the transition region, evidenced by
a blue shift of the c-Si mode, ~c-Si~, from 516 to 520 cm-1. When R
continually increases beyond the transition region, the SRO of the
a-Si:H network, the c-Si volume fraction, and the c-Si peak frequency
are somewhat saturated. However, the structure would be continually
modified at a higher hydrogen dilution ratio. Other techniques such as
TEM will assist in clarifying the further structural changes.

Finally, when the films were deposited with a high growth rate \~20 Å/s
using a high SiH~4~ flow rate, the FWHM of the TO mode for a-Si network
can be as narrow as 58.5 cm^-1^ and the peak frequency as low as 472
cm^-1^. A column-type growth occurs in such conditions, which implies
that both the ~TO~ and the FWHM are sensitive to the column structures.
Calculations are needed to find how the phonon density of states changes
with such structures.

**References**

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II\. Micro-Raman Studies of Mixed-Phase Hydrogenated Silicon Solar Cells

## II.1 INTRODUCTION

A significant light-induced enhancement in the open-circuit voltage
(V~oc~) of mixed-phase silicon solar cells was observed \[1, 2\]. Since
the average grain size is in the order of a few nano-meters in the
mixed-phase material, the magnitude of the variation in V~oc~ (V~oc~)
depends on the intrinsic (*i*) layer thickness, the deposition
temperature, the initial V~oc~, and the light soaking intensity. Under
intense light soaking, a V~oc~ as large as 150 mV has been observed.
Subsequent annealing of the solar cells substantially restored the
original V~oc~ values. *In-situ* photoluminescence (PL) spectral studies
found that the PL intensity and peak-energy position associated with the
amorphous component of the mixed-phase material increase upon light
soaking. We proposed that a reduction of nanocrystalline volume fraction
(X~c~) or size is responsible for the V~o~ enhancement \[1, 2\].
However, recent studies by conventional Raman and X-ray diffraction
spectroscopy found no observable changes before and after light soaking,
prompting us to carry out further investigations.

Raman spectroscopy is a sensitive tool that provides valuable structural
information about hydrogenated amorphous silicon (a-Si:H) materials. For
the mixed-phase material, the X~c~ is low, and the distribution of the
grains might be non-uniform. Using conventional Raman to detect the
existence of nano-grains is difficult. However, micro-Raman is an
effective tool for detecting nano-grains and providing their
distribution information. In this paper, we present the micro-Raman date
on the mixed-phase solar cells, which correlated well the structural
non-uniformity to the physical model of light-induced increase of V~oc~
observed in the mixed-phase solar cells.

## II.2 eXPERIMENTAL 

Single-junction *n i p* solar cells were deposited by United Solar onto
4 cm 4 cm stainless steel (ss) substrates using a conventional *rf* glow
discharge technique. Indium-tin-oxide (ITO) dotes of area 0.25 cm^2^
were deposited on the *p* layer as the top contact. The deposition
parameters were adjusted to arrive at the amorphous-to-nanocrystalline
transition region \[2\].

The Raman spectra were collected under ambient conditions using the
514.5 nm line of an argon-ion laser. The penetration depth of the
514.5-nm light is \~60 nm for a-Si:H and is larger for nc-Si:H. The
spectral resolution was approximately 0.5 cm^-1^ in a frequency range of
100 to 1000 cm^-1^. The frequency of the Raman lines were calibrated
using the TO mode of a (111) c-Si wafer at 520 cm^-1^. A low laser power
was used to avoid possibility of laser-induced crystallization. The
incident beam was approximately three microns in diameter. At first,
four Raman scans were taken on the bottom quadrant of the dots. The
locations of the scans are shown in Fig. 1(a). Since non-uniformity
appears to be more pronounced for the cells with V~oc~ \~ 0.75 V, twenty
scans were made in steps of 500 µm in the same bottom quarter of sample
13678-42, as indicated in Fig. 1(b). The laser beam was moved to a new
spot on the ITO dot. This allowed us to scan the samples under the same
initial conditions for each scan and thereby avoid any residual effects
(such as light induced metastabel effects) of laser radiation at the
previously irradiated spot.

![](media/image26.wmf){width="1.9027777777777777in" height="2.125in"}
(a) (b)

![](media/image27.wmf){width="1.9027777777777777in" height="2.125in"}

**Figure 1.** Schematics of (a) 4 scans for all samples, and (b) 20
scans for 13678-42. The diameter of the dot is \~3 µm.

## II.3 RESULTS AND DISCUSSION

![](media/image28.wmf){width="3.438888888888889in" height="3.5in"}

A cell\'s V~oc~ can be conveniently used to determine if it is in the
amorphous or mixed phase \[1, 2\]. Cells with high, low, and medium
V~oc~, corresponding to a-Si:H, nc-Si:H and mixed-phase Si solar cells
were studied. Table I lists the V~oc~ values and the Raman features for
four scans on each cell. The four scans for sample 13687-32 with V~OC~ =
0.981 V are the same and exhibit the typical amorphous signature, as
seen in Fig. 2. The spectrum can be decomposed into four broad Gaussian
distributions: one centered at 480 cm^-1^ with a FWHM of 67 cm^-1^
representing the TO mode, one at \~ 140 cm^-1^ with FWHM of 100 cm^-1^
representing the TA mode, one around \~330 cm^-1^ with FWHM of 117
cm^-1^ representing the LA mode, and one around \~ 440 cm^-1^ with FWHM
of

FWHM of 100 cm^-1^ representing the LO mode. The four components in the
Raman spectrum are typical for high quality a-Si:H material. Figure 3
shows the Raman TO mode for sample 13678-13 with a V~OC~= 0.674 V. All
Raman spectra indicate a nc-Si:H material with a sharp c-Si TO mode
around 516 cm^-1^ appears. We decompose the Raman spectrum to calculate
X~c~ using

X~c~ =(I~c~ + I~gb~)/\[I~c~ + I~gb~ +y(L)I~a~\], (1)

where I~c~ + I~gb~ and I~a~ are integrated intensities of the peaks
associated with c-Si, grain boudary \[5\], or intermediate order \[6\],
and a-Si:H components, respectively. y(L) decreases from 1 to 0.4 when
the grain size (L) grows from 3 nm to 30 nm. We obtain a volume fraction
of 35%-56%.

**Table 1.** Sample V~OC~ with Raman results

When the mixed-phase cells with V~OC~ \~ 0.767 V and \~ 0.748 V were
scanned by Raman, interesting results were observed as shown in Figs. 4
(a) and 4 (b). The four scans in Fig. 4(a) show typical amorphous
features. In Fig. 4(b), three scans show the amorphous signature, but
one scan clearly exhibits a sharp Lorentzian peak around 516 cm^-1^. The
fact that only one of the eight spectra shows nanocrystallinity
indicates that the material is very non-uniform in the micrometer scale.
One may question if the cell in Fig. 4 (a) also contains nc-Si, because
the V~OC~ \~ 0.767 V corresponds to a mixed-phase structure. The sample
in Fig. 4 (a) was further scanned in 20 systematic steps of 500 m
illustrated in Figure 1 (b). Since the beam size is only 3 m in
diameter, there was no possible overlap between adjacent scans.

![](media/image29.wmf){width="2.8944444444444444in" height="3.3in"}

![](media/image30.wmf){width="3.0in" height="2.8055555555555554in"}

**Figure 4.** Raman shifts of mixed-phase cells with (a) V~OC~ = 0.767 V
and (b) V~OC~ = 0.748 V.

The 20 Raman spectra shown in Fig. 5 exhibit typical Gaussian lineshapes
centered around 480 cm^-1^. Several spectra show a hint of a
nanocrystalline shoulder around 512 cm^-1^. Of the entire set of
spectra, only one spectrum shows definite nanocrystallinity. This
confirms the structural non-uniformity for both the mixed-phase cells
with V~OC~ \~ 0.767 V and \~ 0.748 V. Since only a small portion of the
entire cell was scanned, the overall nanocrystalline volume fraction is
difficult to determine. For the c-Si peak on 13678-43 and 13678-42, the
frequency was at 516.7 cm^-1^ and 514 cm^-1^, respectively. When the
c-Si grain size is as small as a few nm, the momentum conservation will
be relaxed and Raman active modes will not be limited to the center of
the Brillouin zone. Thus, the frequency could shift from 520 cm^-1^ to
512 cm^-1^ with decreasing grain size from 10 to 3 nm \[5, 7, 8\].
Grains with sizes smaller than 3 nm are beyond the Raman detection limit
and their scattering would contribute to the a-Si mode. It is clear that
the small c-Si grains are not uniformly distributed in the a-Si:H matrix
but are clustered in a micrometer-size spot.

In the two-diode model \[9\], a mixed-phase solar cell is equivalent to
two diodes connected in parallel. One diode has the characteristics of
an a-Si:H cell with V~oc~ around 1 V, and the other

![](media/image32.wmf){width="1.375in" height="4.5in"}

![](media/image33.wmf){width="1.375in" height="4.5in"}

![](media/image36.wmf){width="1.375in" height="4.5in"}

Raman Shift (cm^-1^)

**Figure 5.** Raman Shift of 20 scans on a mixed phase cell 13678-42.

one has the characteristics of an nc-Si:H cell with a V~oc~ below 0.5 V.
It is assumed in the model that a certain area of the cell contains more
nanocrystallites than the others. In this case, the lateral transport
between the a-Si:H phase and the nc-Si:H phase becomes negligible. The
non-uniform distribution of nanocrystallites obtained by micro-Raman
measurements in this study provides a direct supporting evidence for the
assumption.

## II.4 SUMMARY

We used micro-Raman spectroscopy to study the degree and dimension of
the structural uniformity of mixed-phase silicon solar cells. Low-laser
power was used to avoid possible thermal effects. For a cell with V~oc~
= 0.981 V, four scans on different locations of the cell show typical
broad Gaussian lineshapes around 480 cm^-1^, indicative of an amorphous
silicon structure. All scans of the cell with V~oc~ = 0.674 V show a
sharp Lorentzian peak around 516 cm^-1^, with an estimated 40%
nanocrystalline volume fraction. For mixed-phase cells with V~oc~ =
0.748 and 0.767 V, one of twenty and one of four spectra, respectively,
exhibited both a broad Gaussian lineshape around 480 cm^-1^ and a sharp
peak around 516 cm^-1^. This suggests that a large fraction of a-Si:H
component but small fraction of nc-Si:H grains within the mm dot of the
mixed phase cells. The spatial distribution of the nc-Si:H grains is
clustered in a micrometer spot. The light-induced V~oc~ enhancement has
been observed in such mixed-phase cells. The structural non-uniformity
observed in this micro-Raman study provides direct supporting evidence
for the teo-diode model recently proposed for explaining the
light-induced V~oc~ enhancement in mixed-phase solar cells.

## 

## REFERENCES

\[1\] K. Lord, B. Yan, J. Yang, and S. Guha, *Appl. Phys. Lett.* **79**,
(2001) 3800.

\[2\] J. Yang, K. Lord, B. Yan, A. Banerjee, S. Guha, D. Han, and K.
Wang, *Mater. Res. Soc. Symp. Proc*. **715** (2002) 601.

\[3\] J.S. Lannin, in Semiconductors and Semimetals, Vol 21, part B,
chapter 6, edited by J. L. Pankove, (Academic Press, Inc., London,
1984).

\[4\] G. Yue, J. D. Lorentzen, J. Lin, Q. Wang and Daxing Han, *Appl.
Phys. Lett.* **75** (1999) 492.

\[5\] S. Veprek, Z. Iqbal, and F. A. Sarott, Philos. Mag. B **45**,
(1982) 137.

\[6\] D. V. Tsu, B. S. Chao, S. R. Ovshinsky, S. Guha, and J. Yang,
Appl. Phys. Lett. **71**, (2003) 1317.

\[7\] S. Veprek, F. A. Sarott, and Z. Iqbal, Phys. Rev. B **36**, (1987)
3344.

\[8\] Y. He, C.Y. Yin, G. Cheng, L. Wang, X. Liu, and G.Y. Hu, J. Appl.
Phys, **75** (1994) 797.

\[9\] B. Yan, G. Yue, J. Yang, K. Lord, and S. Guha, submitted to Appl.
Phys. Lett. (2003).

**III. Toward understanding the degradation without light soaking in
hot-wire a-Si:H thin films and solar cells**

**III.1 introduction**

The advantages of the hot wire chemical vapor deposition (HWCVD) are the
simplicity in gas decomposition, the absence of ion bombardment, and the
high growth rate compared with the standard plasma enhanced (PE) CVD.
High performance hydrogenated amorphous silicon (a-Si:H) based solar
cells and thin-film transistors have been made using HWCVD \[1, 2\]. In
addition, several new physical phenomena were also observed in the HW
a-Si:H thin films \[3-6\]. For instance, the NMR studies showed a new
hydrogen distribution in HW films compared to PECVD films \[3\], the
internal fraction studies showed HW films have a much lower friction
value than glass and are very close to the value of crystalline silicon
\[4\]. In the high growth-rate films, the findings of the confinement
effect on spin-spin interactions in nanogas systems \[5\] and
nanovoid-related large red shift of photoluminescence peak energy have
been reported \[6\].

The degradations in HW solar cells and films were different than in the
PECVD ones. The features of light soaking effects in HW a-Si:H films
were distinguished from those in conventional PECVD films. The
photoconductivity (PC) decreased only slightly and the dark conductivity
increased in the films deposited at a substrate temperature ≥ 320ºC
\[7-9\]. The increase of the density of the metastable dangling bonds
(D^o^) was found to be the same as that in PECVD films. On the other
hand, it has been reported that in addition to the light-induced
degradation the HW solar cells showed a few percent decrease in
efficiency without exposure to the light \[10\]. Most degradation in
fill factor has shown in the air but not in the vacuum or in liquid
nitrogen environment. Those results suggest a possible post-oxidation in
the devices. The degradation without light soaking is a more serious
problem in the HWCVD samples than that in the PECVD samples. In this
work, we studied in detail the variations in electronic properties of
intrinsic HW films in terms of gas adsorption and thermal annealing in
order to understand the phenomena of degradation in the dark.

**III.2 Experimental DETAILS**

Intrinsic HW a-Si:H films with low initial dark conductivity and large
E~a~ ≥ 0.85 eV were selected as listed in Table I. One standard PECVD
film, A7119, was used for comparison. Group 2 was a pair of films with
and without a thin SiN~x~ passivition layer for post-oxidation study.
Group 3 was one sample that was cut into three pieces for gas adsorption
and thermal annealing studies. Corning-1737-glass substrate was used for
all the samples. The coplanar metal contacts were first deposited on the
glass. The details of the HWCVD system are published elsewhere \[1\].

Table 1 Sample information

+------------+-----------------+-------------+------------------------+
| Group      | Sample          | T~s~ (°C)   | Growth Rate            |
|            |                 |             |                        |
|            | ID              |             | (Å/s)                  |
+------------+-----------------+-------------+------------------------+
| 1          | T773            | 360         | 10                     |
+------------+-----------------+-------------+------------------------+
|            | T779            | 360         | 10                     |
+------------+-----------------+-------------+------------------------+
|            | T837            | 400         | 12                     |
+------------+-----------------+-------------+------------------------+
| 2          | H1441           | 360         | 14                     |
+------------+-----------------+-------------+------------------------+
|            | H1442           | 360         | 14                     |
+------------+-----------------+-------------+------------------------+
| 3          | H1540a          | 320         | 12                     |
+------------+-----------------+-------------+------------------------+
|            | H1540b          | 320         | 12                     |
+------------+-----------------+-------------+------------------------+
|            | H1540c          | 320         | 12                     |
+------------+-----------------+-------------+------------------------+
|            | A7119^a^        | 200         | 1                      |
+------------+-----------------+-------------+------------------------+

^a^ PECVD sample

The measurements were carried out in the following experiments: (a) Dark
conductivity measurement at 300 K in environments of: vacuum (10^-4^
Torr), dry-oxygen gas, and air. Thermal annealing cycles were followed
after storage in the above environments. (b) Conductivity
activation-energy measurement with sample temperature ranging from 300
to 420 K. Before taken data, the sample was heated to 420 K for 30 min
to clean up the film surface adsorption. Then, the E~a~ was deduced from
the equation ~d~ = ~o~exp(-E~a~/kT). (c) Hot-probe measurement is a
simple way to distinguish between n- and p-type semiconductors using a
hot-probe such as a heated soldering iron, a cold-probe and a standard
multimeter. Based on the thermal-electric effect of semiconductors, when
applying the probes to n- or p-type materials one obtains a current in
opposite directions. Finally, PC temperature dependence was studied for
two typical HW and one PECVD film.

## III.3 Result and discussion

For group 1 HW samples, the PC decreased only slightly and the dark
conductivity increased upon light soaking \[7-9\]. We used hot probe to
exam the conductive type of the intrinsic samples. We found that high
T~s~ HW samples appeared slightly p-type initially and changed to n-type
after defect creation. Figure 1 shows the results of the hot-probe
measurement for sample T837. One can see that the sample at the initial
state was slightly p-type, and then changed to slightly n-type after the
metastable defect creation. In corresponding to its E~a~ value, the
Fermi level moved up from 0.92 eV to 0.82 eV below the conduction band
edge.

  ---------------------------------------------------- -------------------------------------------------------------------
  ![](media/image37.wmf){width="2.727777777777778in"   Figure 1. The sign of current for sample T837 measured by hot
  height="2.9in"}                                      probe. The current was positive at the initial state where E~a~ was
                                                       as large as 0.92 eV. After the metastable defect creation at State
                                                       B2 (light soaked using 200 mW/cm^2^ white light for 2 hrs) and B6
                                                       (light soaked using 200 mW/cm^2^ white light for 6 hrs), the
                                                       current became negative, and the E~a~ decreased to 0.82 eV.

  ---------------------------------------------------- -------------------------------------------------------------------

This change is in an opposite direction from standard PECVD film (A7119)
in that the E~a~ increased from 0.73 eV to 0.77 eV after the metastable
defect creation. The changes in the activation energy and the PC were
partially recovered after annealing at 160 ^°^C for an hour. The
activation energy recovered to \~0.86 eV instead of 0.92 eV. So, there
is some non-reversible degradation. In order to exam the effect of
post-oxidation on the position of the Fermi level, we studied a pair of
film with and without the SiN~x~ passivition layer. We assumed that a
thin layer of SiN~x~ would prevent adsorption from air. The dark current
as a function of temperature and annealing time are shown in Fig. 2.

![](media/image38.wmf){width="2.8472222222222223in"
height="2.6166666666666667in"}

![](media/image39.wmf){width="2.9305555555555554in"
height="2.6166666666666667in"}

> Figure 2. (a) Current-temperature dependence and (b) current as a
> function of annealing time at 420 K, for samples of H1441 and H1442
> with- and without SiN~x~ passivition layer. The samples were mounted
> in the vacuum and in pure oxygen or nitrogen, respectively.

Figure 2 shows the results of (a) conductivity temperature dependence in
various gas environments, and (b) the conductivity as a function of
annealing time at 420 K. As we expected, the H1442 without a SiN~x~
passivition layer showed conductivity changes in both measurements in
regard to the gases of O~2~ and N~2~; but, there were very little
changes for the sample H1441 with SiN~x~ layer. The 10 times difference
in the value of the current between H1441 and H1442 can be attributed to
the interface states between SiN~x~ and a-Si:H. This interface effect
dominates the conductivity measurement on H1441.

For solar cells, we have reported a 9.8% initial efficiency that became
9.2% after a few weeks of being stored in the dark. The degradation
mainly came from the fill factor (FF). In order to check whether the
degradation took place in the intrinsic layer, we cut sample H1540 into
three pieces of H1540a, H1540b and H1541c and studied the current
changes in the vacuum (x), pure oxygen () and the air (∆),
simultaneously. Figures 3(a) and 3(b) show the changes of the dark
current at 300 K and the corresponding activation energy, respectively.
The crosses in Fig. 3(a) show that the conductivity in the vacuum of
10^-4^ torr did not change much up to a month. However, the conductivity
increased about an order of magnitude after following thermal cycles
heated up to 420 K. The change in conductivity corresponded to E~a~
changes from 0.81 to 0.78 eV as shown in Fig. 3(b). The open circles
show interesting changes of the conductivity in pure O~2~: the initial
dark current has the same value as in the vacuum but increased and then
decreased. Whereas, the current value at 300 K increased about an order
of magnitude after temperature increased to 420 K; again the changes in
conductivity corresponded to E~a~ changes as shown in Fig. 3(b).
Finally, the open triangles show the conductivity measurement in the
air: the dark current increased about 60 times more than that in the
vacuum within one day and continuously increased in the air for 54 days.
The current decreased an order of magnitude after E~a~ measurements in
the air. We believe that this is due to de-adsorption of humidity at
elevated temperatures.

![](media/image40.wmf){width="2.8333333333333335in"
height="2.861111111111111in"}

![](media/image41.wmf){width="2.9027777777777777in"
height="2.861111111111111in"}

Figure 3 Degradation experiments: (a) the dark current at 300 K for
samples of H1540a in vacuum (), H1540b in pure O~2~ (), and H1540c in
the air (∆). The data following the arrows indicate the current after
E~a~ measurements. (b) Several measurements of E~a~ obtained from those
three samples after the current measurements at 300 K for about a month.
The data following the arrow indicate the measurements were done in
vacuum for H1540c.

As shown in Figs 3(a) and 3(b), the final values of both the current at
300 K and their activation energy E~a~ tended to reach the same value.
The above results suggest that H~2~O from the air has a stronger effect
than the pure O~2~. The source of the degradation in the dark can be
related to trace oxygen incorporation at 300 K in the air or at 420 K in
a 10^-4^ torr vacuum (we can not exclude the fact that the structural
relaxation may also play a role) that results in an upward movement of
the Fermi-level position. Consequently, the hole-mobility-lifetime
product decreases and then the FF decreases.

  ----------------------------------------------------- -----------------------------------------------------
  ![](media/image42.wmf){width="3.0416666666666665in"   ![](media/image43.wmf){width="3.0416666666666665in"
  height="2.9305555555555554in"}                        height="2.9305555555555554in"}

                                                        
  ----------------------------------------------------- -----------------------------------------------------

**Figure 4** PC temperature dependence for (a) HW film T779 and (b)
PE-CVD film A7119.

All of the above results showed that the changes in activation energy,
in other words, the Fermi-level position were crucial to understanding
the degradation in the dark. Since the gap states were crucial for the
Fermi-level position, we compared the gap states between HW and PECVD
films in Figure 4 using PC temperature dependence \[12\]. To show the
opposite movement of the Fermi-level positions in these two films, the
dark conductivity curves are also plotted in Figs 4(a) and 4(b). We
found that there are two extra PC enhancement peaks at \~200 K and 250 K
for HW film, and the \~200 K peak disappeared after light soaking. No
matter what those states are, the disappearance of the deep traps is
consistent with the upward movement of the Fermi level. One may also
notice that the PC value does not decrease at \~ room temperature at
State B for the HW film but it does for the PECVD film.

**III.4 Conclusions**

For HW a-Si:H samples, we have observed the decreases in the FF of the
solar cells and the dramatic changes in the dark conductivity for the
i-layer regarding the various gases and thermal cycles. Three effects
could be involved in the changes that are: surface adsorption, oxygen
incorporation and thermal relaxation. We observed that the surface
adsorption of H~2~O in the air resulted in a great rising of the dark
conductivity as shown in Fig. 3(a). The value of activation energy was
slightly decreased after temperature cycles in the vacuum. When the
activation energy decreased from \~0.9 eV to \~0.8 eV, the material
changed from slightly p-type to slightly n-type, as shown in Fig. 1.
This resulted in a decrease of the mobility-lifetime product of holes,
and then a decrease of FF. We suggest that the degradation of both
materials and solar cells may share the same origin. With prolonged
storage time in the air at 300 K or being heated to 420 K in a 10^-4^
torr vacuum, oxygen incorporation resulted in the Fermi-level position
moving upward. Essentially, the differences in both the microstructures
\[3,5,9\] and the gap states (see Fig. 4) in HW- from PE-CVD samples
could be in response to the degradation in the dark.

**references**

1.  Qi Wang, Eugene Iwaniczko, Jeffrey Yang, Kenneth Lord, Subhendu
    Guha, Keda Wang, and Daxing Han, 29th PVSC IEEE Orlando, USA, 2002,
    to be published.

2.  R.E.I. Schropp, B. Stannowski, A.M. Brockhoff, P.A.T.T.van
    Veenendaal and J.K. Rath, Mater.Phys.Mech.1, 73-82, (2000).

3.  Y. Wu, J. Todd. Stephen, Daxing Han, J.M. Rutland, R. A. Crandall,
    and H. Mahan, Phys. Rev. Lett. **77** 2049-2052 (1996).

4.  Jonathan Baugh, Alfred Kleinhammes, Daxing Han, Qi Wang & Yue Wu,
    Science, **294,** 1505-1507 (2001).

5.  Liu, Xiao, B. E. White, Jr., R. O. Pohl, E. Iwanizcko, K. M.
    Jones, A. H. Mahan, B. P. Nelson, R. S. Crandall and S. Veprek,
    Phys. Rev. Lett. **78**, 4418 (1997).

6.  Daxing Han, Guozhen Yue, Keda Wang, Jonathan Baugh, Yue Wu, Yueqin
    Xu, and Qi Wang, Appl.Phys. Lett. **80**, 40-42 (2002).

7.  Daxing Han, Jonathan Baugh, Guozhen Yue, and Qi Wang, Phys. Rev. B
    **62**, 7169-7178 (2000).

8.  Daxing Han, Guozhen Yue, H. Habuchi, Eugene Iwaniczko, Qi Wang, Thin
    Solid Films, **395**, 134-137 (Sept. 2001).

9.  Daxing Han, Guozhen Yue, Qi Wang, and Tatsuo Shimizu, Thin Solid
    Films, **430**, 141-144 (2003).

10. P. Stradins et al, abstract in the 2^nd^ International conference on
    CAT-CVD processing, (2002, CO) to be published in Thin Solid Films,
    (2003).

11. M. Hack and M. Shur, J. Appl. Phys. 58, 997 (1985).

12. Daxing Han and H. Fritzsche, J. Non-Cryst. Solids **59/60**, 397
    (1983).

**IV. Is there a correlation between the median component in Raman
spectrum and the (111) side peak in XRD?**

In order to study the above puzzle, a group of samples with varied
hydrogen dilution ratio was deposited at 250°C on SS substrate using
HW-CVD for both Raman and XRD studies. The sample information and Raman
results are listed in Table I.

##### Table I Sample preparation and Raman results

<table style="width:100%;">
<colgroup>
<col style="width: 12%" />
<col style="width: 11%" />
<col style="width: 13%" />
<col style="width: 8%" />
<col style="width: 6%" />
<col style="width: 9%" />
<col style="width: 8%" />
<col style="width: 9%" />
<col style="width: 9%" />
<col style="width: 11%" />
</colgroup>
<tbody>
<tr class="odd">
<td>Sample ID</td>
<td>R = H2:SiH4</td>
<td><p>Thickness</p>
<p>(Å)</p></td>
<td><p>Xgb</p>
<p>(%)</p></td>
<td><p>Xc</p>
<p>(%)</p></td>
<td><p>Xgb/Xc</p>
<p>(%)</p></td>
<td><p>Peaka</p>
<p>(cm<sup>-1</sup>)</p></td>
<td><p>Peakgb</p>
<p>(cm<sup>-1</sup>)</p></td>
<td><p>Peakc</p>
<p>(cm<sup>-1</sup>)</p></td>
<td><p>FWHMc</p>
<p>(cm<sup>-1</sup>)</p></td>
</tr>
<tr class="even">
<td>H1592*</td>
<td>1.2</td>
<td>11854</td>
<td>27</td>
<td>73</td>
<td>38</td>
<td>480</td>
<td>500</td>
<td>513.5</td>
<td>11</td>
</tr>
<tr class="odd">
<td>H1595</td>
<td>1.2</td>
<td>10875</td>
<td>24</td>
<td>67</td>
<td>35</td>
<td>480</td>
<td>503</td>
<td>514.0</td>
<td>11</td>
</tr>
<tr class="even">
<td>H1593</td>
<td>10.2</td>
<td>11591</td>
<td>26</td>
<td>89</td>
<td>29</td>
<td>480</td>
<td>504</td>
<td>517.7</td>
<td>8</td>
</tr>
<tr class="odd">
<td>H1594</td>
<td>18</td>
<td>N/A</td>
<td>12</td>
<td>93</td>
<td>13</td>
<td>480</td>
<td>504</td>
<td>517.8</td>
<td>8</td>
</tr>
</tbody>
</table>

Note: H1592 was set at a low pressure for 7 min, then corrected to the
right pressure for 45 min on top. H1594 can\'t be measured optically
because it is too milky.

![](media/image44.wmf){width="3.3055555555555554in" height="2.7in"}

![](media/image47.wmf){width="3.25in" height="2.8in"}

Fig. 1 Raman TO model for H1592, H1595, H1593 and H1594.

For each sample, 3 spots were scanned to check the structural
uniformity. The Raman TO models in three locations of each sample are
shown in Fig. 1. The spectra were deconvoluted into three components:
a-Si:H, c-Si and the third component that we refer to as grain
boundaries or small particles of c-Si. The c-Si volume fraction, the
percentage of the grain boundary, the peak energy and the width of the
c-Si peaks are listed in Table I. We have double-checked the small peak
at 496 cm-1 of H1593. It was always there for the four scans. We don\'t
know about the corresponding structures. It is also unclear if there is
a correlation between the median component in Raman spectrum and the
(111) side peak in XRD.

**V. PL study of growth rate dependence on solar cells made by BPSolar**

Four pairs of BP-Solar films, L2127, L2129, L2136, and L2137 at original
State A and light-soaked state B (50 ^o^C, 1 sun, Na vapor lamp for 600
h) were studied by PL spectroscopy. We have also done the measurements
after in-situ light soaking using a 200 mW/cm^2^ white light at 40 ^o^C
for two hours. The sample conditions and their PL parameters are listed
in table I. The PL energy spectra were measured at the temperature range
of 80 K - 300 K.

Table I The sample preparing conditions and their PL parameters

+--------+----------+---------+-------------------+-------------------+
| Sample | Growth   | Th      | Peak ~main~ at 80 | Slope of I~main~, |
| ID     | rate     | ickness | K,                | T~L~              |
|        |          |         |                   |                   |
|        | (Å/s)    | (µm)    | (eV)              | \(K\)             |
+--------+----------+---------+-------------------+-------------------+
| L2127  | 1        | 1.79    | 1.35              | 24                |
+--------+----------+---------+-------------------+-------------------+
| L2129  | 3        | 1.47    | 1.36              | 27                |
+--------+----------+---------+-------------------+-------------------+
| L2136  | 5        | 1.65    | 1.36              | 28                |
+--------+----------+---------+-------------------+-------------------+
| L2137  | 10       | 1.76    | 1.34              | 40                |
+--------+----------+---------+-------------------+-------------------+

where Peak~main~ and I~main~ are the peak energy position and the
integral intensity of the PL main band.

![](media/image48.wmf){width="3.7in" height="3.4194444444444443in"}

Fig. 1 PL spectra of samples L2127, L2129, L2316 Fig. 2 PL spectra of
sample L2127 with

and L2317 with growth rate of 1, 3, 5 and 10 Å/s growth rate of 1 Å/s at
low temperatures.

at 80 K.![](media/image49.wmf){width="3.6in"
height="3.5027777777777778in"}

The PL spectra at low temperatures (T \< 150 K) are dominated by the
tail to tail radiative transitions that yield the PL main band. The low
energy defect PL band is relatively weak and can be measured at a higher
temperature. Figure 1 shows the PL spectra at 80 K for those four
samples listed in table I. The spectral lineshapes show the same
features as that of typical device-quality intrinsic a-Si:H; whereas,
the peak energy position, Peak~main~, is slightly different. The highest
growth rate film shows the lowest Peak~main~ that can be related to a
broader band tail and/or a narrower optical gap. Figure 2 shows the PL
spectra at 80-160 K for sample L2127. One can see that both the PL
intensity and the Peak~main~, decrease when the temperature increases
due to carrier-themolization in the band tail states. In order to
extract the information of the electronic tail states and the dangling
bond states, all the PL spectra were deconvoluted into two elements: the
main band peaked at \~1.4 eV at 80 K and the defect band peaked at 0.88
eV. The Peak~main~ position and the integral intensity, I~main~, as a
function of temperature at State A for those four films are plotted in
Figs. 3(a) and (b). One can see that for the samples with a growth rate
from 1 to 5 Å/s the temperature dependence of both Peak~main~ and
I~main~ is as normal as that for typical device-quality a-Si:H. However,
for the sample with a growth rate of 10 Å/s both the Peak~main~ and
I~main~ show much weaker temperature dependence. Especially, the PL peak
energy is \"saturated\" at \~ 1.25 eV when T \> 160 K. This indicates
that a deep trapping state exists in this high growth rate sample. The
slope of the I~main~, T~L~, is listed in table I. T~L~= T~o~/ln(~o~) is
related to the width of the band tail, kT~o~, according to the mode of
carrier-themolization in band tail states.

![](media/image50.wmf){width="3.5277777777777777in" height="3.2in"}

![](media/image51.wmf){width="3.5555555555555554in" height="3.2in"}

> Fig. 3 (a) Peak~main~ and (b) I~main~ of samples L2127, L2129, L2316
> and L2317 as a function of temperature

In-situ light soaking for all samples were done. The I~main~ decreased
slightly and the integral defect PL intensity, I~defect~, clearly
increased after light soaking for two hours. Figure 4 shows the
light-induced increase of I~defect~ at 260 K.

![](media/image52.wmf){width="3.5555555555555554in"
height="3.4305555555555554in"}

> Fig. 4 Relative intensity of the I~defect~ at 260 K as a function of
> the sample growth rate before and after in-situ light soaking.

All the data suggest that increasing the deposition rate from 1 to 5 Å/s
has a modest effect on the electronic tail states and a gradual increase
of defect states. The latter, perhaps, will not affect the solar cell
performance. Especially, the features of the samples with a growth rate
of 3 and 5 Å/s are very close to each other. An increase of the growth
rate to 10 Å/s has a dramatic effect in the electronic states. In
additional to the broader tail states and higher density of dangling
bond states, the deep-trapping states may be related to the microvoids.
The material is much more defective compared to the others, that is in
consisting with the results of the microstructural studies (see Don\'s
CSM results, 10, 02, 2002). The total PL intensity decreased and the
relative defect PL intensity increased after light-induced metastable
defect creation.

**VI. Review Article and Publication list in 2003**

## VI.1 Review Article 

We have completed a long article on *Photo- and electro-luminiscence of
a-Si:H and mixed-phase alloys* for a special issue of "Solar Energy
Materials and Solar Cells" **78** (2003) 181-233. This article
summarizes our recent 10-year work of both electoluminescence (EL), the
despersive-transport-controlled recombination, and photoluminescece (PL)
on a-Si:H and mixed-phase alloy.

**VI.2 Publication list in 2003**

1.  Raman study of thin films from amorphous-to-microcrystalline silicon
    prepared by hot-wire chemical vapor deposition, Daxing Han, J. D.
    Lorentzen, J. Weinberg-Wolf, L. E. McNeil, and Qi Wang, J. Appl.
    Phys., **94**, 5 (2003) 2930-2936.

2.  Micro-Raman studies of Mixed-Phase Hydrogenated Silicon Solar
    Cells, J. Owens, D. Han, B. Yan, J. Yang, K. Lord, and S. Guha,
    Amorphous and Heterogeneous Silicon-Based Films, MRS Symposium Proc.
    **762**, edited by J.R. Abelson, G. Ganguly, H. Matsumura, J.
    Robertson, and E. Schiff, A4.5, (2003).

3.  Toward understanding the degradation without light soaking in
    hot-wire a-Si:H thin films and solar cells, Qi Wang, Keda Wang, and
    Daxing Han, MRS Symposium Proc. **762**, edited J.R. Abelson, G.
    Ganguly, H. Matsumura, J. Robertson, and E. Schiff, A7.10, (2003).

4.  Combinatorial approach to thin-film silicon materials and devices,
    Qi Wang, Leandro R. Tessler, Helio Moutinho, Bobby To, John Perkins,
    Daxing Han, and Howard M. Branz, MRS Symposium Proc. **762**, edited
    by J.R. Abelson, G. Ganguly, H. Matsumura, J. Robertson, and E.
    Schiff, A9.1, (2003).

5.  Micro-Raman Measurements of Mixed-Phase Hydrogenated Silicon Solar
    Cells, J. Owens, D. Han, B. Yan, J. Yang, K. Lord, and S. Guha,
    Proc. NCPV review meeting (2003, 3, 21-24, Denver, CO).

6.  Correlation between cell performance and film properties as function
    of growth rate, G. Ganguly, Daxing Han, et al, Proc. NCPV review
    meeting (2003, 3, 21-24, Denver, CO).

7.  Narrow Gap a-SiGe:H Grown by HWCVD, Brent P. Nelson, Daxing Han, M.
    Boshta, B. Alavi, Rubin Braunstein, Don L. Williamson, and Yueqin
    Xu, Proc. NCPV review meeting (2003, 3, 21-24, Denver, CO).

8.  Hydrogen Distribution, Nanostructures and Optical Properties of High
    Deposition Rate Hot-Wire CVD a-Si:H, Jonathan Baugh, Keda Wang,
    Daxing Han, and Yue Wu, Thin Solid Films 430 (2003) 95-99.

9.  Neutral Dangling Bonds May Not Be The Dominant Recombination Centers
    For Photoconductivity In Hot-Wire a-Si:H, Daxing Han, Guozhen Yue,
    Qi Wang, Tatsuo Shimizu, Thin Solid Films 430 (2003) 141-144.

10. Photo- and electro-luminiscence of a-Si:H and mixed-phase alloys,
    Daxing Han and Keda Wang, "Critical Review of Amorphous and
    Microcrystalline Silicon Materials and Solar Cells" a special issue
    of Solar Energy Materials and Solar Cells, 78 (2003) 181-233.

11. Hydrogen structures and the optoelectronic properties in transition
    films from amorphous to microcrystalline silicon prepared by
    hot-wire chemical vapor deposition, Daxing Han, Keda Wang,
    Jessica M. Owens, Lynn Gedvilas, Brent Nelson, Hitoe Habuchi and
    Masako Tanaka, J. Appl. Phys. **93** (2003) 3776.