**Quarterly Status Report**

1 August 2005 to 31 October 2005

NREL Subcontract No. XXL-5-44205-09

Principal Investigator: P. Craig Taylor

Organization: University of Utah

**Innovative Characterization of Amorphous and Thin-Film Silicon for
Improved Module Performance**

Research results during the third quarter of Phase I of NREL Subcontract
XXL-5-44205-09 are reported. During this quarter we have continued
collaborations with United Solar Ovonics Corporation on defects that
contribute to the Staebler-Wronski effect in modules made using a-Si:H
and a-Si~x~Ge~1-x~:H intrinsic layers. Specifically, we have received a
second set of device-quality samples, half of which have been light
soaked. We are starting our NMR experiments on these two samples to
confirm our earlier results on the metastable hydrogen doublet site at
better signal-to-noise ratios. In addition, we have initiated studies of
light soaking at 77 K to determine if the Staebler-Wronski effect
saturates at this temperature or grows continuously as occurs in the
case of the tritium decay in the tritiated samples (as shown below in
this report). In this quarter we have made the greatest progress on our
continuing collaborations with NREL on defects generated in tritiated
a-Si:H.

The most important metastable effect in the prototypical amorphous
semiconductor, hydrogenated amorphous silicon (a-Si:H), is the so-called
silicon dangling bond defect associated with a threefold coordinated
silicon atom. It is this defect that is thought to control the decreases
in the conductivity and photoconductivity that occur due to exposure to
visible light. These changes are collectively known as the
Staebler-Wronski effect \[1\], which is the major factor in determining
the stability of photovoltaic devices and detectors made from a-Si:H.

In this quarterly status report we describe recent results using a novel
method for producing silicon dangling-bond defects by replacing some of
the hydrogen ^1^H with tritium ^3^H. When the tritium, which is
radioactive and decays to ^3^He with a half-life of approximately 12
years, is bonded to silicon its decay leaves behind a silicon dangling
bond and an inert helium atom. \[Tritium decays to helium-3 emitting an
antineutrino and a beta particle with an average energy of 5.7 keV\].
The great advantage of this method for producing silicon dangling bonds
is that one can calculate accurately the number that is produced as a
function of time.

For many years hydrogen diffusion has been invoked to explain the
kinetics for the production and annealing of the Staebler-Wronski effect
\[2-10\]. Although there is recent evidence that hydrogen pairing plays
a role in stabilizing the Staebler-Wronski defects \[11\], the
assumption of significant diffusion of hydrogen at room temperature is
supported only by circumstantial evidence. In addition, there has been
speculation for many years that the Staebler-Wronski effect involves the
creation of far more extensive changes in the local structural order
\[12\] than might be expected from the simple production of a small
number (approximately 10^17^ cm^-3^) of silicon dangling-bond defects.
In this letter we provide direct evidence for the diffusion of hydrogen
to heal the production of a large density (approximately 10^21^ cm^-3^)
of silicon dangling bonds. Some of the diffusion might be essentially
athermal and occur even at very low temperatures. Together these two
results -- the diffusion of hydrogen and the annealing of a large
density of silicon dangling bonds -- greatly enhance our understanding
of the Staebler-Wronski effect, which is the quintessential
metastability in all of the tetrahedrally coordinated amorphous
semiconductors. Our results have a major impact on the two classes of
models \[7, 8\] that have been proposed to explain the Staebler-Wronski
effect because both classes invoke the motion of only a small density of
hydrogen atoms, about one per defect created, to stabilize the defects.

We present results on two samples, deposited by plasma-enhanced chemical
vapor deposition (PECVD) from SiH~4~ and tritium gas mixtures, under
conditions described elsewhere\[13\]. The first sample was grown at a
substrate temperature of 225 ^o^C. This sample is similar to those used
in devices, in which the initial defect densities are the lowest
experimentally obtainable (film thickness of 0.26 μm). The second sample
was grown at 150 ^o^C. This sample, which is representative of a more
defective material (film thickness of 1.5 μm), showed considerable
dihydride bonding (silicon bonded to two Si and two H atoms) as measured
by Fourier transform infrared (FTIR) spectroscopy. Because of the
presence of silicon-tritium bonds, one expects the tritium decay to
accumulate Si dangling bond defects after the samples have been
deposited. The density of these defects should follow approximately the
number of decayed tritium atoms per unit volume. Since our first
measurements were made seven years after deposition, the density of
decayed tritium was about 6x10^20^ cm^-3^ or a few at. %. Surprisingly,
both electron spin resonance (ESR) measurements of the neutral silicon
dangling bond density and photothermal deflection spectroscopy (PDS)
measurements of the absorption of both neutral and charged defects
yielded densities that were lower by about 3 orders of magnitude (5 x
10^17^ cm^-3^ and 4 x 10^17^cm^-3^ for films deposited at 150 ^o^C and
225 ^o^C, respectively). In addition, most of these defects annealed at
temperatures between about 150 and 200 ^o^C with kinetics similar to
that observed for the optically induced defects created in the
Staebler-Wronski effect. The residual spin densities in the samples
after annealing at the highest temperatures, near the deposition
temperatures, were 9 x 10^16^ cm^-3^ and 2 x 10^16^ cm^-3^ for samples
deposited at 150 and 225 ^o^C, respectively. For the film deposited at
225 ^o^C, Fig. 1 shows the fraction of metastable defects that remains
after 30 min at several stepwise annealing temperatures, T~A~ (open
circles). The defect densities are normalized to their saturated values
before annealing. A typical annealing curve for the optically induced
defects in non-tritiated a-Si:H is also shown in Fig. 1 by dark circles.
This film, deposited at the University of Chicago by PECVD at 230 ^o^C
\[14\] was exposed at room temperature to uniformly absorbed cw light of
intensity 4x10^21^ photons-cm^-3^s^-1^ for one week. At the end of this
exposure, the density of defects was approximately 4x10^16^ cm^-3^. The
similarity of the annealing curves for tritiated and non-tritiated
samples suggests similar distributions of defect annealing energies.

![](media/image1.wmf){width="2.875in" height="2.4493055555555556in"}

> **Fig. 1**. Fraction of defects, N~S~, remaining after stepwise
> annealing for 30 min at each temperature T~A~ . Open circles represent
> the annealing of a tritiated sample of a-Si:H made at a substrate
> temperature of 225 ^o^C and stored 7 years at room temperature. Filled
> circles represent the annealing of light-induced defects in
> (non-tritiated) a-Si:H created by a one-week exposure to light.

After the final annealing, new defects begin to accumulate in these
tritiated samples in the dark. The accumulation kinetics depends on the
temperature at which the samples are stored. Figures 2 and 3 show the
number of newly created defects as functions of time after the final
annealing for the samples grown at 225 ^o^C and 150 ^o^C, respectively.
Data are shown for the accumulation of defects at 25, 77, 300, and 353 K
(the last temperature is only shown for the sample grown at 150 ^o^C).
PDS measurements were performed for the samples kept at room
temperature. Because the sample deposited at 225 ^0^C was only 0.26 μm
thick, a considerable portion of the spin density in the annealed state
is likely due to surface or interface defects \[8, 15\]. Therefore, the
annealed spin density was subtracted from the total spin density to get
the concentrations N~S~ of newly created defects shown in Figs. 2 and 3.
The large scatter in the 300 K data at short times in Fig. 2 is
indicative of the inherent error in the measurements.

> ![](media/image2.wmf){width="2.5in"
> height="2.2465277777777777in"}**Fig. 2.** Kinetics of the accumulation
> of new defects, ΔN~S~, after annealing at 180 ^o^C in a sample of
> tritiated a-Si:H deposited at 225 ^o^C. Open circles and triangles
> represent ESR measurements on the sample kept at 77 K and 25 K,
> respectively. Solid symbols indicate the same sample kept at room
> temperature, with the solid circles and diamonds representing ESR and
> PDS measurements, respectively. The dashed curve is a linear fit to
> the 77 K data, and the solid curve show the calculated number of
> tritium decays in the sample after the annealing. The measurement
> error in ΔN~S~ is approximately ±2 x 10^16^ cm^-3^.
>
> ![](media/image3.wmf){width="2.5416666666666665in"
> height="2.2291666666666665in"}**Fig. 3.** Kinetics of the accumulation
> of new defects after annealing at 150 ^o^C in a sample deposited at
> 150 ^o^C. The symbols represent same quantities as those described in
> the caption to Fig.2. The solid line is the calculated tritium decay
> (1.7x10^16^ cm^-3^/h), and the dashed line is a linear fit to the 77K
> data. Large and small solid circles denote two separate experiments at
> room temperature. Solid triangles represent accumulation of defects at
> 80^o^C followed by room temperature ESR measurements. The relative
> error in ΔN~S~ is about 5x10^15^ cm^­3^.

Two distinct features are apparent from the data in Figs. 2 and 3. At
short times the growth is nearly linear, within experimental error, but
much slower than the estimated production of silicon dangling bonds.
Therefore the first general feature is that the initial growth rate is
essentially independent of temperature (athermal) over the entire range
of temperatures studied. At the higher temperatures (300 K and above)
the defects accumulate more slowly at longer times and eventually
saturate, and the saturated densities increase with decreasing
temperature consistent with the data shown in Fig. 1. At 300 K the
growth roughly follows a stretched exponential rise to a saturation that
occurs after approximately 10^4^ hours in the sample grown at 225 ^o^C
and at shorter times (and higher densities) in the sample grown at 150
^o^C. This behavior is very similar to that which is observed under
optical excitation. In contrast, when the sample is kept at 77 K, there
is no sign of saturation even after 10^4^ hours. Therefore, the second
general feature is a saturation level, which is strongly temperature
dependent. If these saturation levels are thermally activated, they
suggest an activation energy of approximately 0.3 eV.

Because the nuclear reaction for tritium decay also creates an energetic
electron (beta particle) and a much less energetic ^3^He, we must first
rule out any annealing (or generating) effects of these two particles.
The ^3^He has a recoil energy of only a few eV, which is unlikely to
have much effect in either annealing or generating defects. The
electrons are emitted with an average energy of approximately 6 keV, and
the average distance they travel through the film is approximately 0.6
μm. During this process each electron produces about 10^3^ electron-hole
pairs, which can in principle generate additional defects; however, this
production rate of electron-hole pairs is equivalent to illumination by
only about 10^-6^ W/cm^2^ of visible light, which implies the production
of about one defect for 50 electrons \[16\]. This production rate is too
small to be of any consequence.

It is conceivable that the energetic electrons locally heat up the
lattice in such a way as to anneal defects, but calculations suggest
that the probability is low. Nonetheless, we tested this defect
annealing hypothesis by exposing the film grown at 225 ^o^C, after
saturation of the defect density, to a 7 keV electron flux in an
electron microscope. The average incident current density was 10^-8^
A/cm^2^ and the exposure time was 1.5 h. This flux is about 200 times
larger than the effective beta particle flux, which bombards any small
volume element in the film (\~ 10^-11^ A/cm^2^). Thus, the 1.5 h e-beam
exposure is equivalent to 10 days of intrinsic beta particle exposure.
In the saturated steady state, the annealing rate is equal to the defect
creation rate of about 10^16^ cm^-3^ h^-1^. If the beta particles were
responsible for annealing of defects, we would see a decrease of the
saturated defect density by about 10^18^ cm^-3^ after the e-beam
exposure, but we observed no measurable change. We conclude that the
emission of beta particles (energetic electrons) during the decay of
tritium has no significant influence on the production or annealing of
defects in tritiated a-Si:H.

In addition, it is conceivable that there is a large density of
positively and negatively charged defects that are unseen by ESR. Figs.
2 and 3 show clearly that this is not the case at room temperature; here
the defect densities measured by PDS (optical absorption), which is
believed to measure both neutral and charged defects, match those
measured by ESR. At 25 and 77K, the growth of the ESR signal is
remarkably similar to that seen at short times at room temperature,
which strongly suggests that charged defects are not important at any
temperature.

It remains to explain the two general features of the growth curves --
an annealing of defects at short times that is essentially athermal, and
an annealing at long times (saturation) that is strongly temperature
dependent. Because we have tested the saturation over time scales
approaching seven years, this feature requires a significant reservoir
of hydrogen sites on the order of an atomic percent. In a typical sample
of a-Si:H, such as our tritiated one grown at 225 ^o^C, there is
approximately 10 at. % hydrogen that ^1^H nuclear magnetic resonance
(NMR) measurements show \[17, 18\] exists in three primary forms: \~ 7
at. % bonded to silicon in small clustered environments, \~ 3 at. %
bonded to silicon at isolated (randomly distributed) sites, and \~ 1 at.
% interstitial molecular hydrogen (H~2~), of which approximately half is
HT (hydrogen-tritium) and one quarter is T~2~ (tritium-tritium) in the
tritiated sample.

The most obvious candidate to provide the mobile hydrogen necessary for
annealing \[19\] is the decay of hydrogen--tritium (HT) or T~2~
molecules, which leaves behind an unbonded H (or T). It is unlikely that
molecular hydrogen is the only source given the very different
temperature dependences for the short-time and long-time annealing
processes and the need for a very large reservoir of potential hydrogen
sites. A second possible candidate is the replacement of two near
neighbor bonds in the clustered phase \[20\] (e.g., Si-H and Si-T) with
silicon-silicon bonds on formation of a silicon dangling bond by tritium
decay. This reconstruction eliminates the newly created Si dangling bond
and at the same time, produces an extra H (or T) atom that can easily
diffuse and passivate another defect nearby.

The low rates of defect production in the linear regime show there is
some tritium decay that does not produce defects: some athermal process
circumvents this defect production. Either the diffusion of at least
some of the hydrogen is athermal or the replacement of some of the
silicon-hydrogen bonds with silicon-silicon bonds, as discussed above,
is athermal. The latter appears to be the more likely possibility if the
lattice relaxation is strong enough at the clustered H sites. This
explanation requires that the H produced in the same process has no ESR
signal. With this mechanism, the metastable defects are created
essentially at the isolated tritium sites, which represent only a
fraction of the total tritium sites. Thus, tritium decay produces both
Si dangling-bond defects at isolated tritium sites and "mobile" hydrogen
at clustered tritium and hydrogen sites. In the saturation regime at
higher temperatures, the hydrogen released at clustered sites can
thermally diffuse to passivate defects. Saturation is possible when the
annealing rate due to diffusion of interstitial hydrogen is equal to the
defect production rate.

What are the implications of these results for understanding the
Staebler-Wronski effect in a-Si:H? The similarity of the annealing of
the defects induced by tritium decay and the optically induced defects
at room temperature (Fig.1) suggests that the annealing in both cases is
likely to be governed by the same mechanism. The annealing of
light-induced defects has been extensively studied and is characterized
by a wide distribution of annealing activation energies \[8, 21\]. There
is much circumstantial evidence that the annealing of optically induced
defects is mediated by mobile hydrogen \[6, 22\]. Both classes of model
for the Staebler-Wronski effect suggest \[7, 8\] that defect annealing
is mediated by a metastable configuration that is created together with
the defects and can reverse the creation process during anneal. In both
models, however, only a small density (10^17^ cm^-3^ in device quality
films) of such metastable species is created after prolonged optical
excitation. In the tritiated samples, however, the tritium bonded to
silicon is converted into inert ^3^He and therefore plays no role in the
formation of metastable configurations capable of annealing Si dangling
bond defects. As discussed above, the metastability must be mediated by
hydrogen from a large reservoir of *already existing* metastable sites
(a few at. %). As no new defects are created by the tritium decay at
these metastable sites, an isolated Si-H could not fill this role, but
molecular hydrogen and clustered hydrogen bonded to silicon could. We
therefore conclude that the annealing of optically induced defects that
are created in the Staebler-Wronski effect at room temperature and
elevated temperatures might also take place from a large, pre-existing
reservoir of annealing species rather than from stabilization mechanisms
of a small density of defects that are created by the light.

Finally, we address the growth of defects created by tritium decay and
those created by optical excitation at temperatures below 300 K. The
optically induced defect densities are difficult to measure at low
temperatures by ESR \[23\] or by the constant photocurrent method \[12,
14\], but neither of these experiments has observed saturation at
temperatures much below 300 K. Unless the presence of optical excitation
plays a role in the saturation process, one can speculate that the
behavior at low temperatures is similar to that observed in the
tritiated samples.

In summary, we have shown direct evidence for the diffusion of hydrogen
to heal the production of a large density (approximately 10^21^ cm^-3^)
of silicon dangling bonds. Some of the network healing mechanisms, which
follow tritium decay, are essentially athermal and occur even at very
low temperatures.

In addition to our collaborative studies on tritiated samples with NREL,
we are continuing our studies on hydrogen doublets with United Solar
Ovonics. We hope eventually to perform measurements on a-Ge:H in
collaboration with United Solar Ovonics, but because of the poorer
quality of the a-Ge:H samples, we have started by investigating the
quality using electron spin resonance and photothermal deflection
spectroscopy. In spite of efforts to optimize sample quality our
experiments so far have shown that the defect densities are too high to
perform the light soaking experiments. When these experiments are
successful, we will plan the NMR measurements on large-area samples. In
addition, we are continuing to examine additional ways of obtaining
samples purposely designed by United Solar Ovonics Corporation to have
much greater degradation than the device quality films. Finally, we are
irradiating device quality samples of a-Si:H at 77 K for long periods of
time to test for saturation.

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