Low substrate temperature amorphous and microcrystalline silicon
films deposited by Radio-Frequency and Hot-Wire Chemical Vapor
Deposition
P. Alpuim, V. Chu
Instituto de Engenharia de Sistemas e Computadores (INESC),
Rua Alves Redol, 9, 1000 Lisboa, Portugal
J. P. Conde
Department of Materials Engineering, Instituto Superior Técnico,
1049-001 Lisboa, Portugal
PACS:
73.61.Tm, 81.05.Ys, 71.23.Cq, 61.43.Dq
ABSTRACT
The effect of hydrogen dilution on the optical, transport and structural properties of
amorphous and microcrystalline silicon thin films deposited by hot-wire chemical
vapor deposition (HW) and radiofrequency plasma-enhanced chemical vapor
deposition (RF) using substrate temperatures (Tsub) of 100 ºC and 25 ºC is reported.
Microcrystalline silicon (c-Si:H) is obtained using HW with a large crystalline
fraction and a crystallite size of 30 nm for hydrogen dilutions above 85%
independently of Tsub. The deposition of c-Si:H by RF, with a crystallite size of 8
nm, requires increasing the hydrogen dilution and shows decreasing crystalline
fraction as Tsub is decreased. The photoconductivity, defect density and structure factor
of the amorphous silicon films (a-Si:H) are strongly improved by the use of hydrogen
dilution in the Tsub range studied. a-Si:H films with a photoconductivity-to-dark
conductivity ratio above 105, a deep defect density below 1017 cm-3, an Urbach energy
below 60 meV and a structure factor below 0.1 were obtained for RF films down to 25
ºC (at growth rates 0.1-0.4 Å/s) and for HW films down to 100ºC (at growth rates
10 Å/s), using the appropriate hydrogen dilution. In the low Tsub range studied, the
growth mechanism, film properties and the amorphous to microcrystalline silicon
transition depend on the flux of atomic hydrogen available. The properties of the
films are compared to those of samples produced at 175 ºC and 250 ºC in the same
reactors.
2
I. INTRODUCTION
Hydrogenated amorphous silicon (a-Si:H ) is widely used in solar cells and other large
area electronic applications such as active matrix flat panel displays or light sensors in
image scanners.1 This is due to the possibility of depositing this material in large areas
at low costs and to its good optical and electronic properties.1,2 However, because of
metastability in a-Si:H, which deteriorates its properties upon light-exposure or bias
application, as well as its low carrier mobility (1 cm2V-1s-1) and low doping
efficiency,1,3 microcrystalline silicon (c-Si:H) is currently being investigated as a
strong candidate for a next generation material for many of these device
applications.4,5
Presently, most large area applications use either glass or stainless steel sheets as
substrates. However, new applications which require a flexible, lightweight,
unbreakable and economical substrate such as plastic have brought about the need to
develop low temperature growth and processing of amorphous and microcrystalline
silicon.6,7 Because of the much lower glass transition temperature of plastic compared
with glass, the deposition temperature of a-Si:H and c-Si:H must be lowered from
the 250-350 ºC range that is currently used to under 150 ºC,7 while maintaining the
electronic properties.
Device quality a-Si:H and c-Si:H are normally deposited by RF glow discharge from
silane, where the microcrystalline fraction of the films is controlled by hydrogen
dilution of the reactant gas.2,8 The deposition rates are low, both for a-Si:H from
undiluted SiH4 (rd1 Å/s)9 and for c-Si:H (rd0.2 Å/s) because of the higher
hydrogen dilutions (>95%) required for microcrystalline growth.8 A much faster
3
technique for growing device-quality a-Si:H and c-Si:H films is hot-wire chemical
vapor deposition (HW-CVD) which can give deposition rates 5 to 25 times higher
than those obtained with RF-PECVD.4,10 This is believed to be due to the high
efficiency of the hot tungsten filament (Tfil=2500 ºC) in breaking the silane and
hydrogen molecules.10 The large quantity of atomic hydrogen supplied to the surface
of the growing film is responsible for the good electronic quality of films deposited at
substrate temperatures of 150-350 ºC,2,8 as it passivates dangling bonds and
reconstructs stressed Si-Si bonds in the subsurface.10 At hydrogen dilutions equal to or
greater than 90%, the extent of the reconstruction is such that the film becomes
microcrystalline.3
Pioneering studies on a-Si:H have already reported that for deposition near room
temperature, films have a much higher hydrogen content than for those deposited at
200-300 ºC.11,12 This hydrogen was bonded in more defective ways (SiH2, SiH3 or
(SiH2)n) thus giving rise to poor electronic properties such as low photoconductivity
and high sub-gap absorption. However, the aim of those early studies was to optimize
the deposition of a-Si:H rather than investigate extreme substrate temperatures. Since
the practical importance of very low substrate temperatures (<150 ºC) was recognized,
more detailed work has been done. Hishikawa et al.13 studied a-Si:H films deposited
at 80 ºC and 50 ºC by conventional PECVD and stressed the importance of lowering
the deposition rate to give time for surface reactions to occur and thus increase the
quality of the films. Roca i Cabarrocas14 produced dense and ordered intrinsic, ndoped and p-doped films at Tsub=50 ºC by avoiding both plasma and surface
polymerization reactions.
4
Cheung et al.15 deposited films with a compact network structure at 50 ºC by
optimizing the silane partial pressure in the deposition chamber, while keeping rf
power and Tsub constant. They concluded that successful incorporation of hydrogen in
the films depends on several deposition variables (including but not exclusively the
substrate temperature) that can be tuned for good performance at 50 ºC. Srinivasan et
al.16 produced a-Si:H films at 35 ºC with a high monohydride fraction of bonded
hydrogen using helium dilution of the reactant gases. In that work, a direct correlation
between deposition rate and monohydride dominant hydrogen bonding was not found
but rather the necessity of using an intermediate rf power (10W).
All these materials deposited at low substrate temperatures have in common poorer
optoelectronic properties and higher defect densities than their high deposition
temperature (or annealed) counterparts. However, the question remains if these asdeposited low substrate temperature films are high-quality enough to make a device.
The first low-temperature thin film transistors (TFTs) reported were deposited on
glass, at 150 ºC by Feng et al.17, and at 125 ºC by McCorrnick18 et al. Recently, the
first TFTs on plastic were fabricated, at 125 ºC on polycarbonate by Gates6 and at 150
ºC on polyimide by Gleskova et al.7. The layers for the devices produced at 125 ºC
were rf reactive magnetron sputtered6,18 while those produced at 150 ºC used RFPECVD.7,17 In the PVD technique the kinetic energy of the sputtered atoms from the
target (10-20 eV for the faster atoms6) is transferred to the surface of the growing film
thus supplying it with some thermal activation energy.18 CVD techniques, allowing
conformal step coverage, merit further investigation of their properties in the lowest
deposition temperature range. The only report of a complete TFT on plastic until now
(including n+ contact layer and showing device performance data) is the RF-PECVD
device of ref. 7.
5
In the present work, a-Si:H and c-Si:H thin films were deposited on glass substrates
using HW-CVD and RF-PECVD with substrate temperatures during deposition of 25
ºC and 100 ºC. At each temperature, a series of films grown with increasing hydrogen
dilution of silane was made: 0 to 95% H2 dilution was used in the HW case, and 0 to
99% in the RF case. The properties of these films were compared with those of
samples produced at 175 ºC and 250 ºC in the same reactors. The objective of this
paper is to study the structural and optoelectronic properties of amorphous and
microcrystalline films deposited at room temperature and at 100 ºC.
II. EXPERIMENTAL PROCEDURES
A. Film Preparation
The films were deposited by HW-CVD and RF-PECVD in an ultrahigh-vacuumquality system with two different reactors and a common load lock. Each chamber has
its own pumping system (2 Balzers TPU520) and the base pressure was better than 5 
10-8 Torr. The sample was clamped to the grounded upper electrode which was,
except for the room temperature depositions, heated to the temperature of deposition,
Tsub. The deposition temperature was monitored using a thermocouple embedded in a
copper block in contact with the substrate. The gases, silane and hydrogen, entered the
deposition chamber through a ring-shaped showerhead of 12.5 cm diameter located
approximately 8 cm below the substrate holder.
For HW deposition, a single tungsten filament of 0.5 mm diameter and approximately
7 cm length was placed 5 cm from the substrate and was resistively heated with a DC
power supply. The filament temperature was measured with an optical pyrometer (Tfil
 2500 ºC) and the pressure was kept constant at 20 mTorr.4,10 The thickness of the
6
films was between ~0.2 m and ~0.8 m. The c-Si:H films were, in general, thinner
than the amorphous ones. For the RF deposition, the inter-electrode distance was 3
cm, the density of RF power used was 50 mW/cm2 (in all but the room temperature
depositions where it was 100 mW/cm2), and the pressure was 100 mTorr. Both in the
HW and RF depositions the sum of the fluxes of the gases was kept at around 10
sccm, except for the higher dilutions where it was necessary to increase the fluxes so
that the SiH4 flux was not less than 0.5 sccm, which was the lower limit for the silane
mass flow controller.
B. Film Characterization.
Films were deposited on two different substrates simultaneously: Corning 7059 glass
and double-side polished 100 Si. The first was used for optical transmission, parallel
transport measurements, deep defect density measurements by the Constant
Photocurrent Method, x-ray diffraction and Raman scattering measurements; the
second was used for Fourier transform infrared (FTIR) spectroscopy.
The film thicknesses were measured with a profilometer. The thicknesses obtained
were in good agreement with those calculated from interference fringes of
near-infrared optical transmission.19
The optical bandgap of the films, Eopt, was determined by measuring the transmission
of above-bandgap light through the film.20 Eopt was extracted from the fit of the data
to the Tauc equation E  BE  Eopt , where B is a constant,  is the absorption
coefficient and E the photon energy. The dark conductivity d was measured between
110 ºC and room temperature on coplanar Cr contacts, 6 mm long , 1 mm apart and
approximately 1000 Å thick. The activation energy Ea was calculated from
7
 d   0 exp  Ea / k BT  . The steady-state photoconductivity ph was measured as a
function of generation rate. The light from a 250 W tungsten-halogen lamp was
filtered with a bandpass filter at a wavelength which gives an approximately uniform
carrier generation throughout the thickness of the film. The generation rate, G, was
calculated from the response of a calibrated silicon photodiode located next to the
sample. In this work, ph refers to the photoconductivity at a carrier generation rate of
1021 cm-3 s-1.
The constant photocurrent method, CPM, was used to measure the subgap
absorption.21 The deep defect density Ns of the amorphous films was calculated from
Ns=CCPM   (1.2 eV), where CCPM =1016 cm-2.22 The CPM spectra were normalized
by setting the extrapolation of the Urbach tail equal to the absorption from the
transmission measurements at h=Eopt + 0.1 eV (ref. 23) or, for some samples, by
adjusting to the absorption spectrum from photothermal deflection spectroscopy, PDS,
measurements.
The hydrogen content, CH, and the microstructure factor, R, were determined using
infrared spectroscopy. CH was calculated from the density of silicon atoms, NSi=5 
1022 cm-3, and the density of bonded hydrogen atoms, NH (CH =NH/ NH+NSi). NH was
calculated from the integrated absorption coefficient of the Si-H wagging modes
located around 630 cm-1, using N H  As 
s
  
d where As =1.6  1019 cm-2

(ref.24). As was derived for a-Si:H but the same constant is valid for microcrystalline
films as well. 25 R was calculated from the deconvolution of the stretching band into
two peaks, one centered at approximately 2000 cm-1 (I2000) and the other centered
around 2100 cm-1 (I2100), R =I2100/( I2000+ I2100).26
8
Raman spectra were measured in the backscattering geometry using a Raman
microprobe. The 514.5 nm (2.41 eV) laser radiation was obtained from an Ar+ laser.
The power of the incident beam was set below 50 mW to avoid thermally induced
crystallization. For microcrystalline films, the Raman spectrum around the crystalline
silicon transverse optical (TO) peak27 was deconvoluted into their integrated
crystalline Gaussian peak, Ic (~ 520 cm-1), amorphous Gaussian peak, Ia (~ 480
cm-1), and intermediate Gaussian peak, Im (~ 510 cm-1).28,29 The crystalline fraction,
Xc, was calculated from Xc =(Ic + Im )/(Ic + Im + Ia).21 The crystallite size, dRaman, was
calculated from d Raman  2
B /   , where  is the shift of the peak for the c-
Si:H sample compared to that of c-Si, and B=2.0 cm-1nm2.30 For amorphous films, the
bond angle deviation  was related to the full width at half maximum (FWHM)  of
the TO peak centered at 480 cm-1 using the expression: =/6-2.5.31
X-ray diffraction peaks were measured with a Siemens D-5000 x-ray diffractometer
using the Cu K1 line (=1.54056 Å). The samples were measured at grazing
incidence (0.5  and 1) using substrate holder rotation (15 rpm). The crystallite size
dX-ray was calculated from the Scherrer formula d X  ray  k / B cos B  , where k0.9, 
is the wavelength of the X-ray radiation, B is the FWHM of the peaks (in units of 2 )
and B is the angular position of the peak.
III. RESULTS
Figure 1 shows the deposition rate, rd, as a function of the hydrogen dilution of the
silane in the reactive gas mixture (FH2/( FH2+FSiH4) , where FH2 is the flow of hydrogen
and FSiH4 is the flow of silane). In general, the deposition rates are more than one order
of magnitude higher in HW than in RF. Typical values for rd are ~20 Å/s for HW
9
a-Si:H films and ~1 Å/s for RF a-Si:H films. For c-Si:H films rd~2 Å/s for HW and
~0.2 Å/s for RF. For both HW and RF, rd is independent of Tsub in the temperature
range studied (25 ºC< Tsub<250 ºC). The deposition rate increases with increasing
silane concentration in the gas mixture.
As the hydrogen dilution is increased (and, consequently, the silane fraction in the gas
phase is decreased) the films undergo an amorphous to microcrystalline transition. In
the case of RF-deposited films, the hydrogen dilution necessary for this transition to
occur increases with decreasing substrate temperature. While for Tsub=175 ºC the 95%
hydrogen dilution film is amorphous and the 97% is microcrystalline (crystalline
fraction, Xc=66%) (fig. 2 (a)), for Tsub=100 ºC the 97.5% hydrogen dilution film shows
only Xc=6% and 98% hydrogen dilution is necessary to obtain a microcrystalline film
(Xc=61% and crystallite size dX-ray=7.6 nm) (fig. 2 (b)). For Tsub=25 ºC the 98%
hydrogen dilution film is still amorphous and the 99% hydrogen dilution film still
shows a significant amorphous fraction (Xc=42%) (fig. 2 (c)). In the case of HW films
the hydrogen dilution at which the amorphous-to-microcrystalline transition occurs is
independent of Tsub. Figure 3 shows that all the 80% hydrogen dilution films are
amorphous while all 90% hydrogen dilution films are microcrystalline. The crystalline
fraction for the 90% hydrogen dilution films decreases with Tsub: Xc=84% at 175 ºC;
Xc=74% at 100 ºC; and Xc=43% at 25 ºC. The crystallite size is approximately
constant for Tsub =175 ºC and 100 ºC (dX-ray is respectively 37 nm and 39 nm)
decreasing to dX-ray =25 nm for Tsub=25 ºC (figure 4). Although the crystalline fraction
increases further when 95% hydrogen dilution is used (Xc=91%, 87% and 89% for
Tsub=175 ºC, 100 ºC and 25 ºC, respectively) the Raman spectrum suggests a small
decrease (10%) in the crystallite size. The 85% hydrogen dilution samples show in
10
general a transitional character with both significant amorphous and microcrystalline
fractions present.
Figure 5 shows the room-temperature dark conductivity of HW and RF samples as a
function of hydrogen dilution. The d values cluster between 10-11 and 10-10 -1cm-1
for amorphous films. For HW-deposited films the transition from amorphous to
microcrystalline structure is gradual and independent of the substrate temperature
between 25 ºC< Tsub<250 ºC: up to 80% hydrogen dilution, the films show a d
characteristic of a-Si:H films; at 90% hydrogen dilution and above, d is characteristic
of c-Si:H (of the order of 10-5 cm-1); at 85% hydrogen dilution they show an
intermediate value. RF samples show a very different behavior: the transition is abrupt
(within 1-2 % of hydrogen dilution) and Tsub-dependent. At 175 ºC, d characteristic
of c-Si:H is obtained at 97% hydrogen dilution; at 100 ºC, 98% hydrogen dilution is
required, and, at 25 ºC, 99%, in good agreement with the Raman results (figs. 2 and
3). This result indicates that the parameter range to deposit c-Si:H by RF is severely
reduced upon decrease of Tsub, contrary to HW.
Figure 6 shows Ea plotted as a function of hydrogen dilution. A clear difference
between HW and RF amorphous samples can be observed: the values of Ea for HW
samples scatter between 0.8 and 0.9 eV up to hydrogen dilutions of 80% and then
continuously decrease to 0.55 eV at 90% hydrogen dilution. The RF deposited
samples have values of Ea ~0.95 eV until the amorphous-to-microcrystalline transition
occurs. At this point the value of Ea abruptly falls to 0.55 eV. In this region, both
HW and RF samples show some n-type character, the Ea values becoming lower than
that of intrinsic crystalline silicon.
11
Figure 7 shows the photoconductivity ph of HW and RF films. At low hydrogen
dilutions, ph depends strongly on Tsub, in both HW and RF samples. In RF samples,
ph of samples deposited at Tsub=25 ºC is always more than 2 orders of magnitude
lower (~10-9 -1cm-1 at 50% hydrogen dilution and ~10-7 -1cm-1 at 98%) than ph of
samples deposited at Tsub=100 ºC (~10-7 -1cm-1 at 50% hydrogen dilution and ~10-5
-1cm-1 at 98%). Values of ph for amorphous RF films deposited at Tsub=175 ºC are
well above (between ~1.5 and ~2.5 orders of magnitude) those of the same hydrogen
dilution deposited at Tsub=100 ºC and are less dependent on hydrogen dilution (at 175
ºC, ph increases 1 order of magnitude from 0 to 90% hydrogen dilution while at 100
ºC, it increases more than 2 orders of magnitude). In RF, ph reaches a maximum at
hydrogen dilution 90% (dependent on Tsub), then decreases right before the
microcrystalline transition at 95% hydrogen dilution for Tsub=175 ºC and at 97.5%
hydrogen dilution for Tsub=100 ºC suggesting the formation of a more defective, less
compact material. HW films behave in a very different way: ph increases steadily
with increasing hydrogen dilution until it reaches a maximum value after the
amorphous-to-microcrystalline transition. This maximum is reached at the same
hydrogen dilution level (90%) for all Tsub175 ºC. At this hydrogen dilution, all the
HW c-Si:H films show roughly similar values of ph, independently of Tsub (2.410-5
-1cm-1 at 25 ºCph .1.110-4 -1cm-1 at 175 ºC).
Photo-to-dark conductivity ratios, ph/d, or photosensitivity, are shown in figure 8.
In HW, the maximum photosensitivity is obtained at 50-80% hydrogen dilution
independently of Tsub. A decrease of the maximum ph/d is observed in HW films
12
when Tsub is decreased: (ph/d)max~106 for Tsub=220 ºC, ~5104 for Tsub=175 ºC and
100 ºC, and ~103 for Tsub=25 ºC. c-Si:H films, at 90% hydrogen dilution, show
ph/d ~1. For RF samples, ph/d keeps increasing with increasing hydrogen dilution,
reaching its maximum (~90% hydrogen dilution for Tsub=175 ºC, 96% for Tsub=100
ºC, and 98% for Tsub=25 ºC) just before the amorphous-to-microcrystalline transition.
This maximum results from a drop in the value of d (fig. 5) without a similar drop in
the value of ph (fig. 6). After that transition, ph/d abruptly falls to low values (1-10)
within a span of 1-2% of hydrogen dilution. The highest values of ph/d fall between
105 and 106 independently of Tsub. The range of hydrogen dilution for which
ph/d104 is significantly reduced with decreasing Tsub: for Tsub=175 ºC it is 95%,
for Tsub=100 ºC it is between 90 and 97.5%, and for Tsub=25 ºC it is ~98%.
Figure 9 shows the hydrogen content, in atomic per cent, as a function of hydrogen
dilution. In the 25-175 ºC substrate temperature range, the hydrogen content
(CH~15%) in amorphous HW films is weakly dependent on the deposition
temperature and it decreases with hydrogen dilution from CH~18% at no H2 dilution to
CH~15% at 80% H2 dilution. HW a-Si:H samples deposited at 220 ºC reveal
significantly lower H content (CH~7.5%) independently of hydrogen dilution. Parallel
to the trend observed in the conductivity data, the hydrogen content of all HW films
converges to approximately CH~2-6%, after the amorphous to microcrystalline
transition, at 90% hydrogen. At this hydrogen dilution, the H-content of the Tsub=25
ºC c-Si:H film is CH~6% while the 220 ºC film has CH~2.6%. The RF a-Si:H films at
low hydrogen dilution (50%) have the same hydrogen content as HW ones (15%CH
20%). The hydrogen content increases to a maximum right before the amorphous to
microcrystalline transition (at Tsub=175 ºC, CH~23% at 95% hydrogen dilution, at
13
Tsub=100 ºC, CH~28% at 95% hydrogen dilution and at Tsub=25 ºC, CH~30% at 98%).
This agrees with the previous suggestion of the formation of a more porous material
before this transition occurs2 and with the trend observed in the photoconductivity
values already described in an earlier paragraph (see fig. 7). The hydrogen content of
the RF c-Si:H samples is strongly dependent on deposition temperature, increasing
with decreasing Tsub. The minimum CH for Tsub=100 ºC is CH~12% and for Tsub=25 ºC
it is 25%, suggesting the presence of porous intergrain regions or of a low crystalline
fraction.
In figure 10 the structure factor R is plotted as a function of hydrogen dilution. A high
content of monohydride bonded hydrogen (low R) is associated with low density of
electronic traps in a-Si:H. c-Si:H, although having much lower hydrogen content
than amorphous films (fig.9), typically have higher R-values due to the existence of
many terminating Si-H2 units in its grain boundaries, which behave like internal
surfaces or voids. At 0% dilution, the HW a-Si:H films show R values between ~0.25
(Tsub=220 ºC) and 0.75 (Tsub=25 ºC). The value of R decreases until a H2 dilution of
60-80% where R0.1 for Tsub=175 ºC and 100 ºC and R=0.27 for Tsub=25 ºC. An
abrupt increase of R occurs for hydrogen dilution above 80%. The lowest R factors for
RF a-Si:H films deposited at 100 ºC (R=0.01 at 50% hydrogen dilution) and 25 ºC
(R=0.06 at 90% hydrogen dilution) are significantly lower than those for HW a-Si:H
deposited at the same Tsub .
Figure 11 shows the absorption coefficient , measured by CPM, plotted against
photon energy of HW and RF films deposited from undiluted silane. Figure 12 shows
the CPM spectrum which gives the lowest subgap absorption at each substrate
temperature. The best films were considered to be the films with the sharpest Urbach
14
tails and the minimum subgap absorption, estimated from the absorption at 1.2eV. In
general, amorphous samples by HW give broader curves, with higher Urbach tails
than corresponding samples by RF. At 0% hydrogen dilution, the Tsub=220 ºC, 175 ºC
and 100 ºC samples by HW had Eu=57meV, 81meV and 80meV, respectively. The 25
ºC sample was too defective to allow a CPM measurement. Values of Eu of the
undiluted RF films at 250 ºC and 175 ºC were 50meV and 58 meV, respectively, and
the RF sample deposited at Tsub=100 ºC had Eu=85 meV. While at high substrate
temperatures (220 ºC in HW, 250 ºC in RF) the best a-Si:H were obtained without
hydrogen dilution, the 100 ºC amorphous samples required a dilution of 60% in HW
and 95% in RF to approach these results: Eu=59 meV and 1.2=7.8 cm-1 in HW, Eu=50
meV and 1.2=3.0 cm-1 in RF. The best a-Si:H films at Tsub=25 ºC were obtained with
80% hydrogen dilution in HW and 98% in RF, with Eu=80 meV, 1.2=55 cm-1 for the
HW film and Eu=67 meV, 1.2=11 cm-1 for the RF sample. 50% hydrogen dilution
allowed the deposition of the best films at Tsub=175 ºC, both in HW and RF. In most
of the low-Tsub a-Si:H films by HW, values of subgap absorption at 1.2 eV vary
between 10 and 100 cm-1 (1017 cm-2<Ns<1018 cm-2) while the corresponding RF values
vary between 1 and 10 cm-1 (1016 cm-2<Ns<1017 cm-2).
III. DISCUSSION
Table I summarizes the results of undiluted and selected a-Si:H and c-Si:H samples
deposited with hydrogen dilution at Tsub=25 ºC, 100 ºC and 175 ºC by HW and RF.
Also shown for comparison are results for high- Tsub device-quality materials.
Atomic hydrogen plays an important role in the growth process of amorphous and
microcrystalline silicon.2,8,32 When the growth surface is exposed to atomic hydrogen,
15
there can be abstraction of hydrogen bonded to silicon (with consequent creation of
dangling bonds), breakage of weak silicon-silicon bonds,
chemical etching by
forming silane, formation of Si-H bonds with the surface dangling bonds, and
diffusion into the film.33,34 These bond modifications and long-range atomic motion
on the surface promote the formation of a relaxed structure.35,36 In this way, the
presence of atomic hydrogen can compensate a decrease in substrate temperature. This
effect is clearly observed in both RF and HW a-Si:H films where, when Tsub is
decreased, an increasingly higher hydrogen dilution is necessary to achieve high
photoconductivity (fig.7). Atomic hydrogen stabilizes the structure, breaking weak SiSi bonds and promoting cross-linking. The relaxation of the structure induced by these
chemical reactions can achieve a level at which crystallization occurs.35,37 One
distinguishing characteristic of HW is the high rate of atomic hydrogen produced by
H2 or silane dissociation on the hot tungsten filament at Tfil 1900 ºC.36,38,39 The high
atomic hydrogen concentration in HW can be invoked to explain the insensitivity to
Tsub of the hydrogen dilution at which the amorphous to microcrystalline transition
occurs (figs. 3 and 5).
The amorphous to microcrystalline transition shows different characteristics in RF and
HW. While an abrupt transition (within an interval of 1-2% of hydrogen dilution) is
observed in the case of RF films, a more gradual transition (between 80 and 90% of
hydrogen dilution) is observed in the case of HW films (figs. 5 and 6). For the RF
films there is a narrow interval of hydrogen dilution between the range where typical
amorphous and the range where typical microcrystalline films are obtained. Films in
this interval are characterized by an increase in hydrogen content (fig. 9) and of the Rfactor (fig. 10) and a decrease in d (fig. 5). In addition, these films (for Tsub=175 ºC,
16
95% hydrogen dilution, for Tsub=100 ºC, 96% hydrogen dilution and, for Tsub=25 ºC,
98% hydrogen dilution) often show a very small (<10%) crystalline fraction and are
those with highest photo-to-dark conductivity ratio (figs. 2 and 8). In addition, these
particular films also show the lowest values of bond angle disorder (6º) among aSi:H films deposited by RF at Tsub175ºC. It is possible that the nucleation of
crystallites in RF is preceded by the formation of a highly porous, protocrystalline film
which includes small crystallites and a continuous network of high quality amorphous
silicon. This is supported by the CPM subgap absorption results of fig.12 which show
that the films with the lowest defect density for each temperature series are those
which occur just prior to the microcrystalline transistion. However, although these
pre-transition films show similar photo-to-dark conductivity ratios (fig. 8) for all
temperatures, the defect density increases for decreasing Tsub. The photoconductivity
(fig. 7) of these films correspondingly decreases but the photo-to-dark conductivity
ratio is compensatated by a much sharper drop in the dark conductivity (fig. 5). The
decrease in photoconductivity results both from an increase in defect density and from
a lowering of EF, which increase the density of recombination centers. In contrast,
amorphous HW films appear always to contain small highly ordered regions
(protocrystallites) invisible by Raman and XRD. With increasing hydrogen dilution,
these regions give origin to crystallites without the need for the formation of a porous
nucleation layer, as in RF. The presence of these regions could explain the higher
Urbach tail and defect density observed in HW a-Si:H samples with respect to RF
samples (Table I). The absence of a need for a nucleation layer is compatible with the
observation of ultra-thin HW microcrystalline films (d100 Å).
17
IV. CONCLUSIONS
a-Si:H can be prepared by RF with good optoelectronic properties at Tsub as low as 25
ºC. Photosensitivities above 105 and subgap absorption below 10 cm-1 were obtained
for all substrate temperatures studied (25 ºC<Tsub<175 ºC) when the appropriate
hydrogen dilution was used. As Tsub is lowered the range of hydrogen dilution
necessary to obtain good transport properties narrows significantly and moves to
higher values. a-Si:H deposited by HW shows a photosensitivity which is strongly
dependent on substrate temperature (decreasing from 106 at Tsub=220 ºC to 103 at
Tsub=25 ºC). For Tsub as low as 100 ºC it is possible to obtain HW a-Si:H films with
Eu< 60 meV and 1.2 below 10 cm-1, but at Tsub=25 ºC the best values obtained for Eu
and 1.2 were 80 meV and 25 cm-1, respectively. For RF, as Tsub is lowered, the range
of hydrogen dilution necessary to obtain c-Si:H narrows and moves to higher values.
At the same time the crystalline fraction decreases strongly with Tsub. In contrast, for
HW, c-Si:H is obtained with high crystalline fractions (Xc> 80%) independently of
the Tsub used. In addition, the range of hydrogen dilution for which c-Si:H is
obtained using HW is independent of Tsub.
The use of hydrogen dilution allows the improvement of optoelectronic properties of
low-Tsub films in comparison with those deposited from undiluted silane. Using the
appropriate hydrogen dilution, RF a-Si:H films deposited with Tsub down to 25 ºC and
HW a-Si:H films deposited with Tsub down to 100 ºC show properties compatible with
their use as active layers in TFTs. HW c-Si:H films show more promise for
application as active layers in TFTs than RF c-Si:H for 25 ºC< Tsub<100 ºC.
18
ACKNOWLEDGEMENTS
The authors thank R. Almeida (IST/INESC) for the use of the FTIR and Raman
equipment and L. Paramés and O. Conde of the Faculty of Sciences of the University
of Lisbon for the use of the X-ray diffraction equipment. This work was supported by
the Fundação para a Ciência e Tecnologia (FCT) through Pluriannual Contracts with
UCES/ICEMS (IST) and INESC and by projects PRAXIS/3/3.1/MMA/1775/95 and
PRAXIS/3/3.1/MMA/1792/95. One of the authors (P.Alpuim) thanks the Department
of Physics of University of Minho for a leave of absence.
19
Table I. Properties of selected a-Si:H and c-Si:H samples prepared by
HW and RF.
rd
Eopt
d
(%)
(Å/s)
(eV)
(-1cm-1)
(-1cm-1) (eV)
S1263
0
25.3
1.81
4.210-12
1.310-9
1.07 0.85
a)
a)
HW
S1265
80
9,1
1.79
6.610-10
7.210-7
0.88 0.69
80
55
25
HW
S1271
90
1,5
2.09
7.910-6
2.810-5
0.57 0.42
c)
12 d)
25
RF
S1287
50
1.0
1.65
5.910-11
1.010-9
0.67 0.99
a)
a)
25
RF
S1286
98
0.11
2.15
3.110-13
1.410-7
1.13 0.82
67
10.8
25
RF
S1319
99
0.09
2.10
8.610-9
5.910-8
0.63 0.40
c)
c)
100
HW
S1122
0
24.7
1.78
1.310-10
4.110-7
0.85 0.73
80
75
100
HW
S1131
60
12.8
1.76
4.410-11
2.110-6
0.87 0.75
59
7.8
100
HW
S1127
90
1.6
1.96
1.710-5
9.510-5
0.57 0.66
c)
11 d)
100
RF
S1080
0
1.19
1.76
4.610-12
1.210-8
1.06 0.83
85
29.5
100
RF
S1117
95
0.31
1.87
1.810-11
5.510-6
1.02 0.79
50
3.04
100
RF
S1175
98
0.12
2.14
3.010-6
1.710-5
0.49 0.31
c)
17 d)
175
HW
S1254
0
25.3
1.80
1.810-10
4.510-7
0.89 0.68
81
31
175
HW
S1253
50
13.5
1.75
3.810-11
8.710-7
0.85 0.76
80
12.8
175
HW
S1258
90
1.27
1.93
5.210-5
1.010-4
0.43 0.66
c)
123 d)
175
RF
S1074
0
0.82
1.75
8.710-11
5.610-6
0.99 0.77
58
4.07
175
RF
S1280
50
0.37
1.71
2.910-10
1.210-5
0.87 0.75
68
4.75
175
RF
S1311
97
0.19
1.93
8.910-5
5.010-5
0.28 0.45
a)
a)
220
HW
S467
0
27.7
1.70
6.510-11
1.710-5
0.89 0.76
57
1.83
Tsub
CVD Sample H2dil
( ºC)
tech.
25
HW
25
ph
Ea,d

Eucpm
1.2
(meV) (cm-1)
20
220
HW
S465
95
1.67
1.75
7.010-4
9.310-4
0.29 0.58
c)
75d)
250
RF
S1073
0
0.93
1.72
1.110-9
1.810-4
0.93 0.80
50
0.5
250
RF
S888
97.5
2.15
2.210-4
1.910-4
0.27
c)
c)
a)
No signal was detected. Either the sample is too defective or the photo-to-dark conductivity ratio is
not large enough (e.g., c-samples) to allow CPM measurement.
b) No homogeneous film was deposited, due to powder formation in the gas phase.
c)
Microcrystalline sample.
d) Absorption at 1.0eV, 1.0, is quoted for some c-samples.
21
FIGURE CAPTIONS
Figure 1 - Deposition rate, rd, plotted as a function of hydrogen dilution of silane for
films deposited by HW and RF. Since the filament to substrate distance was R=3 cm
for the Tsub=220ºC HW deposition, the corrected values for the standard distance (5
cm, in this article) are also shown, assuming a 1/R dependence of the number of film
precursors that reach a unit substrate area.
Figure 2 - Raman spectra for RF samples deposited at (a) Tsub=175 ºC, (b) Tsub=100
ºC, and (c) Tsub=25 ºC, using different hydrogen dilutions near the amorphous to
microcrystalline transition.
Figure 3 - Raman spectra for HW samples deposited at (a) Tsub=175 ºC, (b) Tsub=100
ºC, and (c) Tsub=25 ºC, using different hydrogen dilutions near the amorphous to
microcrystalline transition.
Figure 4 - X-ray spectra of samples deposited by HW using 90% hydrogen dilution
deposited at (a) Tsub=175 ºC, (b) Tsub=100 ºC, and (c) Tsub=25 ºC.
Figure 5 - Room-temperature dark conductivity d of RF samples (top) and HW
samples (bottom) plotted as a function of hydrogen dilution for different values of
Tsub. The lines are guides to the eye.
Figure 6 - The activation energy of dark conductivity, Ea, plotted as a function of
hydrogen dilution for HW (left side) and RF (right side) samples. Horizontal dashed
lines indicate typical Ea of intrinsic c-Si (lower line) and of intrinsic standard a-Si:H
(upper line). The arrows point out the abrupt amorphous to microcrystalline transition
of RF films at Tsub=25ºC, 100ºC and 175ºC.
22
Figure 7 - Photoconductivity at a generation rate of 1021 cm-3s-1, ph, of RF samples
(top) and HW samples (bottom) plotted as a function of hydrogen dilution for
different values of Tsub. The lines are guides to the eye.
Figure 8 - Photosensitivity ph/d of RF samples (top) and HW samples (bottom)
plotted as a function of hydrogen dilution for different values of Tsub. The lines are
guides to the eye.
Figure 9 - Hydrogen content of HW (left side) and RF (right side) samples plotted as
a function of hydrogen dilution for different values of Tsub. The hydrogen content was
calculated from the integrated 630 cm-1 absorption band in the infrared absorption
spectra. The lines are guides to the eye.
Figure 10 - Structure factor R calculated from the integrated 2000 cm-1 and 2100 cm-1
absorption bands in the infrared absorption spectra of the films, plotted as a function
of hydrogen dilution for HW and RF samples deposited at different temperatures.
The lines for the Tsub=25 ºC and 100 ºC are guides to the eye.
Figure 11 - The optical absorption coefficient  determined by the constant
photocurrent method (CPM) plotted as a function of photon energy for HW (left side)
and RF (right side) samples deposited with 0% hydrogen dilution at different Tsub.
Figure 12 - The optical absorption coefficient  (from CPM measurements) plotted as
a function of photon energy for HW (left side) and RF (right side) a-Si:H samples. At
each Tsub, the sample with the lowest Urbach energy and lowest subgap absorption
was chosen, independently of hydrogen dilution.
23
1
See, for example, R. A. Street, Hydrogenated Amorphous Silicon (Cambridge University Press,
1991).
2
C. C. Tsai, Amorphous Silicon and Related Materials, edited by Hellmut Fritzsche (World Scientific
Publishing Company, 1988) pp.123-147.
3
P. Brogueira,V. Chu, A. C. Ferro, J. P. Conde, J.Vac.Sci.Technol.A 15(6), 2968 (1997).
4
A. H. Mahan, M. Vanacek, A. Poruba, V. Vorlicek, R. S. Crandall, D. L. Williamson, Materials
Research Society Symposium Proceedings 507 (Pittsburg, Pennsylvania: Materials Research Society,
1998), in press.
5
K. Erickson, V. L. Dalal, G. Chumanov, Materials Research Society Symposium Proceedings 467
(Pittsburg, Pennsylvania: Materials Research Society, 1997), p.409.
6
S. M. Gates, Materials Research Society Symposium Proceedings 467 (Pittsburg, Pennsylvania:
Materials Research Society, 1997), p.843.
7
H. Gleskova, S. Wagner, Z. Suo, Materials Research Society Symposium Proceedings 507 (Pittsburg,
Pennsylvania: Materials Research Society, 1998), in press.
8
A. Matsuda, Journal of Non-Crystalline Solids 59-60, 767 (1983).
9
R. A. Street, Mater. Res. Soc. Bull. 17, 70 (1992).
10
J. P. Conde, P. Brogueira, V. Chu, Philosophical Magazine B 76, No. 3, 299 (1997).
11
H. Fritzche, M. Tanielian, C. C. Tsai, P.J.Gaczi, J. Appl. Phys. 50, 3366 (1979).
12
G. Lucovsky, R. J. Nemanich, J.C.Knights, Phys. Rev.B 19, 2064 (1979).
13
Y. Hishikawa, S. Tsuge, N. Nakamura, S. Tsuda, S. Nakano, Y. Kuwano, J. Appl. Phys. 69, 508,
(1991).
14
P. Roca i Cabarrocas, Appl. Phys. Lett. 65, 1674 (1994).
15
M. K. Cheung, M. A. Petrich, J. Appl. Phys. 73, 3237 (1993).
16
E. Srinivasan, D. A. Lloyd, G. N. Parsons, J.Vac.Sci.Technol.A 15(1), 77 (1997).
17
M. S. Feng, C. W. Liang, D. Tseng, J. Electrochem. Soc. 141, 1040 (1994).
24
18
C. S. McCorrnick, C. E. Weber, J. R. Abelson, J.Vac.Sci.Technol.A 15(5), 2770 (1997).
19
A. Goodman, Appl. Opt. 17, 2779 (1978).
20
J. Tauc, R. Grigorivici, A.Vancu, Phys. Status Solidi 15, 627 (1966).
21
M. Vanacek, J. Kocka, J. Strichlik, Z. Kosicek, O. Stika, A. Triska, Sol. Energy Mater 8, 411 (1983).
22
N.Wyrsch, F.Finger,T.J.McMahon, M.Vanacek, J.Non-Cryst. Solids 137-138, 347 (1991).
23
T.Watanabe, K.Azuma, M.Nakatani, T.Shimada, Jpn. J. Appl. Phys. Part 2 29, L1419 (1990).
24
C.J.Fang, K.J.Gruntz, L.Ley, M.Cardona, F.J.Demond, G. Mueller, S.Kalbitzer, J.Non-Cryst. Solids
35-36, 255 (1980).
25
U. Kroll, J. Meier, A. Shah, S. Mikhailov, J. Weber, J. Appl. Phys. 80, 4971 (1996).
26
E.C.Molenbroek, A.H. Mahan, E.J. Johnson, A.C. Gallagher, J. Appl. Phys. 79, 7278 (1996).
27
J. S. Lannin, Semiconductors and Semimetals, Part B 21 (Academic Press, New York, 1984).
28
T. Kaneko, M. Wakagi, K. Onisawa, T. Minemura, Appl. Phys. Lett 64, 1865 (1994).
29
S. Veprek, F. A. Sarott, Z. Iqbal, Phys. Rev.B 36, 3344 (1987).
30
Y. He, C. Ying, G. Cheng, L. Wang, X. Liu, G. Y. Hu, J. Appl. Phys. 75, 797 (1994).
31
Y. Hiroyama, R. Suzuki, Y. Hirano, F. Sato, T. Motooka, Jpn. J. Appl. Phys., Part 1 34, 5515 (1995).
32
I. Shimizu, J.Non-Cryst. Solids 114, 145 (1989).
33
H. Shirai, J.Hanna, I. Shimizu, Jpn. J. Appl. Phys. 30, L881 (1991).
34
H. Shirai, J.Hanna, I. Shimizu, Jpn. J. Appl. Phys. 30, L679 (1991).
35
G. N. Parsons, J. J. Boland, J. C. Tsang, Jpn. J. Appl. Phys. 32, 1943 (1992).
36
H. V. Nguyen, I. An, R. W. Collins, Y. Lu, M. Wagaki, C. R. Wronski, Appl. Phys. Lett..65, 3335
(1994).
37
T. Akasaka, I. Shimizu, J.Non-Cryst. Solids 198-200, 883 (1996).
38
I. Langmuir, G. M. J. Mackay, J. Am. Chem. Soc. 36, 1708 (1914).
39
T. W. Hickmott, J. Appl. Phys. 31, 128 (1960).
25