581
Amorphous an
26. Amorphous and Microcrystalline Silicon
Hydrogenated amorphous silicon (a-Si:H) and microcrystalline silicon (µc-Si:H) are recognized as being
useful materials for constructing devices related to optoelectronics, such as solar cells, thin-film transistors,
etc. [26.1, 2]. Several methods have been proposed
for the preparation of device-grade a-Si:H and µcSi:H. These include: reactive sputtering of a crystalline
silicon target with Ar+H2 plasma [26.3]; mercury- sensitized photochemical vapor deposition (CVD) utilizing
a decomposition reaction of silane (SiH4 ) molecules
with photoexcited Hg (Hg*) [26.4]; a direct photo
CVD method where high-energy photons from a Xeresonance lamp or a low-pressure Hg lamp are used
for the direct excitation of SiH4 molecules to excited
electronic states [26.5, 6]; a hot-wire CVD method for
26.1 Reactions in SiH4 and SiH4 /H2 Plasmas .. 581
26.2 Film Growth on a Surface ..................... 583
26.2.1 Growth of a-Si:H....................... 583
26.2.2 Growth of µc-Si:H ..................... 584
26.3 Defect Density Determination
for a-Si:H and µc-Si:H ......................... 589
26.3.1 Dangling Bond Defects ............... 589
26.3.2 Dangling Bond Defect Density
in µc-Si:H ................................ 590
26.4 Device Applications .............................. 590
26.5 Recent Progress in Material Issues
Related to Thin-Film Silicon Solar Cells .. 591
26.5.1 Controlling Photoinduced
Degradation in a-Si:H................ 591
26.5.2 High Growth Rates
of Device-Grade µc-Si:H ............ 592
26.6 Summary ............................................ 594
References .................................................. 594
decomposing SiH4 by means of catalytic reactions on
a heated metal surface [26.7]; and a plasma-enhanced
CVD method (PECVD). The PECVD method is the
most widely used of these due to its ability to consistently prepare uniform, high-quality materials on
a large-area substrate.
In this chapter, the PECVD method is highlighted,
details regarding the processes used to grow a-Si:H
and µc-Si:H from reactive plasmas are explained, and
the determination reaction (which is used to obtain the
dangling-bond defect density in the films: one of the
most important structural properties that influences device performance) is interpreted in order to obtain clues
about how to control the optoelectronic properties of
those materials for device applications.
26.1 Reactions in SiH4 and SiH4 /H2 Plasmas
The initial event required for the growth of a-Si:H and
µc-Si:H is the decomposition of the source gas material
in SiH4 or SiH4 /H2 glow discharge plasma. Figure 26.1
shows a schematic of the dissociation pathway of SiH4
and H2 , during which the molecules are excited to higher
electronic states due to inelastic collisions with highenergy electrons in the plasma [26.8]. As the electrons
in the plasma usually have a wide variety of energies,
Part C 26
Processes used to grow hydrogenated amorphous
silicon (a-Si:H) and microcrystalline silicon (µcSi:H) from SiH4 and H2 /SiH4 glow discharge plasmas
are reviewed. Differences and similarities between
growth reactions of a-Si:H and µc-Si:H in a plasma
and on a film-growing surface are discussed,
and the process of nucleus formation followed by
epitaxial-like crystal growth is explained as being
unique to µc-Si:H. The application of a reaction
used to determine the dangling-bond defect
density in the resulting a-Si:H and µc-Si:H films
is emphasized, since it can provide clues about
how to improve the optoelectronic properties
of those materials for device applications,
especially thin-film silicon-based solar cells.
Material issues related to the realization of
low-cost and high-efficiency solar cells are
described, and finally recent progress in this area
is reviewed.
582
Part C
Materials for Electronics
SiH4
H2
+
SiHx
Hx+
*
Si
10– 8 s
SiH*
10– 7 s
–
Part C 26.1
SiHx
SiH2
hν
(656 nm)
hν
(288 nm)
H
Si
hν
(414 nm)
SiH3
H*
1.0
SiH
Flux distribution (eV– 1)
H2
H
0.5
0
0
5
10
15
20
Electron energy (eV)
Fig. 26.1 Schematic showing the dissociation of SiH4 and H2 molecules to a variety of chemical species in the plasma
via excited electronic states. The electron energy distribution function in the plasma is also shown
from zero to several tens of electron volts (eV), groundstate electrons of source gas molecules are excited into
their electronic excited states almost simultaneously due
to inelastic collisions with energetic electrons. Excited
electronic states of complicated molecules like SiH4 are
usually dissociating states, from which dissociation occurs spontaneously to SiH3 , SiH2 , SiH, Si, H2 and H,
as shown in Fig. 26.1, depending on the stereochemical structure of the excited state. Hydrogen molecules
are also decomposed to atomic hydrogen. Excitation of
ground-state electron to vacuum-state gives rise to ionization events, generating new electrons and ions, which
maintains the plasma.
Reactive species produced in the plasma also experience secondary reactions, mostly with parent SiH4 and
H2 molecules, as shown in Fig. 26.2, resulting in a steady
state. Reaction rate constants for each reaction are summarized in the literature [26.9]. Steady-state densities of
reactive species are basically determined by the balance
between their rate of generation and their annihilation
rate. Therefore, highly reactive species such as SiH2 ,
SiH and Si (short-lifetime species) have much smaller
densities than SiH3 in the steady-state plasma, although
the generation rates of those species are not very different from that of SiH3 , which shows low reactivity along
with SiH4 and H2 (long-lifetime species).
Steady-state densities of reactive species have
been measured using various gas-phase diagnostic techniques [26.10–16], such as optical emission
Ion exchanging
Ion – Molecule
Neutral – Molecule
Disproportionation
Insertion
Recombination
Abstraction
SiHx+ + SiH4
SiHx + SiH4
SiH + SiH4
Si + SiH4
SiH2 + SiH4
SiH2 + H2
SiH3 + SiH4
H + SiH4
SiHx + SiH4+
SiH3 + SiH3
Si2H5
SiH3 + SiH
Si2H6
SiH4
SiH4 + SiH3
H2 + SiH3
Ion – Radical
less probable
Radical – Radical
Fig. 26.2 Representative secondary reactions of the chemical species produced in the plasma with SiH4 and H2
molecules. Their reaction rate constants are available in
the literature
spectroscopy (OES) [26.10], laser-induced fluorescence (LIF)
[26.12], infrared laser absorption
spectroscopy(IRLAS) [26.13], and ultraviolet light absorption spectroscopy (UVLAS) [26.16]. Figure 26.3
shows the steady state number densities of the chemical species, including both emissive and ionic species,
in the SiH4 and SiH4 H2 plasmas used to prepare devicegrade a-Si:H and µc-Si:H. It is clear from Fig. 26.3 that
the SiH3 radical is the dominant chemical species in the
growth of both a-Si:H and µc-Si:H, although the density
ratio of short-lifetime species to SiH3 changes depending on the conditions used to generate the plasma. For
instance, when high electric power is supplied to the
plasma (high electron density) under low SiH4 flow
Amorphous and Microcrystalline Silicon
SiH2
Si*
SiH3
Si
SiH*
105
SiH
H
106
107
108
109 1010 1011 1012 1013
Number density of species in the steady state plasma (cm–3)
Fig. 26.3 Number densities of chemical species in realistic steadystate plasmas measured or predicted by various diagnostic
techniques
that atomic hydrogen plays an important role in µcSi:H growth [26.17] although SiH3 is the dominant film
precursor for both a-Si:H and µc-Si:H growth [26.18].
26.2 Film Growth on a Surface
26.2.1 Growth of a-Si:H
Upon reaching a film-growing surface, SiH3 radicals begin to diffuse across it. During this diffusion, the SiH3
abstracts bonded hydrogen from the surface, forming
SiH4 and leaving dangling bonds on the surface (this
known as growth site formation). Other SiH3 molecules
then diffuse across the surface to find the site containing the dangling bonds, whereupon Si−Si bond
formation (film growth) occurs, as shown schematically in Fig. 26.4. This surface reaction scheme for
film growth has been proposed on the basis of two
experimental results [26.18].
SiH3
SiH4
SiH3
H
Figure 26.5 shows the general concept of the surface reaction process. Some the flux of SiH3 is reflected
off the surface (the proportion of molecules reflected is
given by the reflection probability). The remaining SiH3
is adsorbed onto the surface and it changes its form as
follows:
1. SiH3 abstracts bonded H from the surface, forming SiH4 , or two of the SiH3 radicals interact on
the surface, forming Si2 H6 (with a recombination
probability γ );
2. Surface-diffusing SiH3 sticks to the site containing the dangling bond, forming Si−Si bond (with
a sticking probability s).
The total loss probability (β) is given by the sum of
recombination probability and the sticking probability
(γ + s), and the reflection probability is therefore 1 − β.
Among the reaction probabilities mentioned above, the
Si
Incident flux 1
SiH3
Reflection (1 –β)
SiH3
Recombination (γ)
SiH4
Si2H6
Surface diffusion
Sticking (s)
a – Si: H
Fig. 26.4 Schematic of the surface growth of a-Si:H
583
Fig. 26.5 General concepts behind the surface reactions of
incoming SiH3 radicals
Part C 26.2
rate conditions, SiH4 is rapidly dissociated and then
depleted, giving rise to reduced probabilities of gasphase reactions of short lifetime species with SiH4
molecules. This leads to an increased contribution from
short-lifetime species to the film growth, which causes
a deterioration of the structural properties of the resulting
films.
The steady-state density of atomic hydrogen (H)
varies widely in the plasma, as shown in Fig. 26.3.
This is mainly due to the change in the hydrogen dilution ratio R (H2 /SiH4 in the source gas material, i. e.,
the density of atomic hydrogen increases with increasing R. Noting the fact that more µc-Si:H is formed with
increasing R at constant electron density in the plasma
and constant substrate (surface) temperature, it is clear
26.2 Film Growth on a Surface
584
Part C
Materials for Electronics
a)
b) Deposition rate (A/s)
1.0
1.5
0.8
Estimated from
step-coverage profiles
0.6
1.0
Part C 26.2
0.4
0.5
0.2
Determined by
a grid method
0
0
100
200
300
400
500
Substrate temperature (°C)
0
0
100
200
300
400
500
Substrate temperature (°C)
Fig. 26.6a,b Loss probability β (a) of SiH3 radicals reaching the film-growing surface, as measured using the grid
method and the step-coverage method, and the deposition rate (corresponding to the sticking probability s for SiH3 ) (b)
as a function of substrate temperature during the formation of a-Si:H
RR power density (W / cm2)
2.0
A
1.5
A
1.3
A
A
A
A
A
1.0
0.8
0.5
A
0.3
0.1
A
A
A
A
0.06
δ 100Å
0.03
Xc 100 %
0.01
SiH4
1
4
7
9
14
19
29
49
H2-dilution ratio (R = H2 / SiH4)
Fig. 26.7 Crystal size (light: δ) and volume % (dark: X c ) of microcrystallites in the resulting films, mapped out on the RF power
density/hydrogen dilution ratio plane
dicted from the film deposition rate as a function of the
substrate temperature, respectively.
The loss probability β, related to the reflectivity (1 − β), is simply dependent on the nature of the
site where SiH3 from the plasma lands. Therefore,
the temperature-independent 1 − β seen in Fig. 26.6
suggests that almost all of the surface sites are covered with bonded H over the whole temperature range
used in the experiment, although a few dangling bonds
are created thermally above 350 ◦ C, as shown by in
situ infrared reflection absorption spectroscopy (IRRAS) [26.19]. On the other hand, SiH3 diffusing across
the surface is easily captured by even just a few dangling bonds, because a mobile species on the surface
can find a specific site at a distant location from the
landing site.
The temperature-independent behavior of 1 − β
and the temperature-dependent behavior of s observed in Figs. 26.6a and 26.6b have been perfectly
reproduced by theoretical simulations based on the
surface-reaction scheme shown in Fig. 26.4, assuming reasonable activation energies for the surface
diffusion of SiH3 , the abstraction reaction of H
with SiH3 , the saturation of the site containing the
dangling bond with SiH3 , and thermal H-removal
processes [26.20].
26.2.2 Growth of µc-Si:H
loss probability has been measured using a grid method
(GM) as well as a step-coverage method (SCM). Figures 26.6a and 26.6b show the loss probabilities obtained
by GM and SCM, as well as the sticking probability pre-
Atomic hydrogen reaching the film-growing surface
plays an important role in the growth of µc-Si:H. This
has been confirmed in the µc-Si:H-formation map draw-
Amorphous and Microcrystalline Silicon
50
Crystallite volume fraction (%)
26.2 Film Growth on a Surface
585
Raman crystallinity (Ic / Ia)
6
Cathode
40
5
R = 19
30
Mesh
4
Anode
Probe
Part C 26.2
9
20
3
4
10
2
0
100
200
300
400
500
Substrate temperature (C)
Fig. 26.8 Volume % (X c ) of microcrystallites in the result-
ing films plotted against the substrate temperature during
film growth
1
0
1
10
100
Ion energy (arb. units)
Fig. 26.9 Raman crystallinity in the resulting µc-Si:H as a function
ing on the RF power density/hydrogen diffusion ratio
plane as shown in Fig. 26.7. As seen in the figure,
large-crystallite µc-Si:H is prepared under high hydrogen dilution conditions, indicating the importance
of atomic H in the growth of µc-Si:H. Figure 26.8
shows the crystalline volume fraction (determined by
the X-ray diffraction peak area) in the resulting µcSi:H as a function of the substrate temperature during
film growth for three different hydrogen dilution ratios [26.17]. The crystalline volume fraction increases
with increasing substrate temperature, reaching a maximum at around 350 ◦ C, and then suddenly drops to zero
above 500 ◦ C. The lack of µc-Si:H formation above
500 ◦ C suggests that surface hydrogen coverage is a
requirement for crystallite formation in the resulting
film [26.17].
Figure 26.9 shows the negative effect of ionic species
impinging on the film-growing surface on the formation
of µc-Si:H, studied using a triode reactor, as shown in
the inset of the figure [26.21]. The Raman crystallinity
Ic /Ia , defined as the ratio of the peak intensity from
the crystalline phase at around 520 cm−1 to that from
the amorphous phase at 480 cm−1 in the Raman scattering spectrum, deteriorates as the energy of the ionic
species impinging on the film-growing surface (controlled by the bias voltage applied to the mesh electrode)
is increased.
The surface reaction behavior of SiH3 reaching the
film-growing surface has also been investigated for µcSi:H growth using GM and SCM, in contrast to a-Si:H
growth [26.18]. Figures 26.10a and 26.10b show the
loss probability (β) and the deposition rate (corresponding to the sticking probability s) of SiH3 reaching the
of the energy of the ions impinging on the film-growing surface
(which is controlled by the application of the bias voltage to the
mesh electrode during film growth)
a) β
1.0
0.8
µ c – Si: H
0.6
0.4
0.2
0
a – Si: H
0
100
200
300
400
500
Substrate temperature (°C)
b) Deposition rate (A / s)
1.2
1.0
0.8
0.6
a – Si: H
0.4
µ c – Si: H
0.2
0
0
100
200
300
400
500
Substrate temperature (°C)
Fig. 26.10 Loss probability β and deposition rate as a
function of substrate temperature for µc-Si:H growth in
comparison to those for a-Si:H growth
film-growing surface as a function of the substrate temperature for the growth of both µc-Si:H and a-Si:H.
586
Part C
Materials for Electronics
Based on these experimental results, the following
properties of the formation process of µc-Si:H can be
specified:
SiH3
H
Enhanced surface
diffusion
Local heating
Sufficient surface
H coverage
Part C 26.2
Fig. 26.11 Surface diffusion model for µc-Si:H formation. The large
spheres and small spheres represent Si and H, respectively
SiH4
SiH4
Growth Models for µc-Si:H
In an attempt to explain the specific phenomena observed during the formation of µc-Si:H, three models
have been proposed:
SiH3
H
Etching
Etching
H-permeation
Amorphous phase
Re-growth
Crystalline phase
Fig. 26.12 Etching model for µc-Si:H formation
SiH3
H
Growth zone
Growth
H-permeation
1. The film precursor is SiH3 , the same as in the case
of a-Si:H growth;
2. Atomic hydrogen reaching the film-growing surface
plays an important role in the formation of µc-Si:H;
3. The film becomes amorphous when the substrate
temperature is higher than 500 ◦ C;
4. High-energy ions impinging on the surface result in
crystallinity deterioration;
5. The surface loss probability shows a temperature
dependence whereas the sticking probability does
not show a temperature dependence in the case of
µc-Si:H growth.
Structural relaxation
Fig. 26.13 Chemical annealing model for µc-Si:H formation
Unlike in the case of a-Si:H growth, the loss probability
β shows a significant dependence on substrate temperature and the sticking probability s shows no temperature
dependence in the case of µc-Si:H growth.
1. The surface diffusion model [26.17];
2. The etching model [26.22];
3. The chemical annealing model [26.23].
The surface diffusion model is depicted schematically in Fig. 26.11. Here, a high atomic H flux from the
plasma results in full bonded hydrogen surface coverage and also local heating through hydrogen exchange
reactions on the film-growing surface. These two actions enhance the surface diffusion of film precursors
(SiH3 ). As a consequence, the SiH3 adsorbed on the
surface can find energetically favorable sites, leading to
the formation of an atomically ordered structure (nucleus formation). After the formation of the nucleus,
epitaxial-like crystal growth takes place with enhanced
surface diffusion of SiH3 [26.17, 24].
The etching model has been proposed due to the
observation that the rate of film growth decreases
with increasing hydrogen dilution ratio R. The concept behind the etching model is shown schematically
in Fig. 26.12. Atomic H reaching the film-growing surface breaks Si−Si bonds, preferentially weak bonds
involved in the amorphous network structure, leading
to the removal of Si atoms weakly bonded to other Si
atoms. This site is replaced with a new film precursor
SiH3 , creating a rigid, strong Si−Si bond, which gives
rise to an ordered structure [26.22, 24].
The chemical annealing model has been proposed
in order to explain the observation that crystal formation is observed during hydrogen plasma treatment;
Amorphous and Microcrystalline Silicon
Formation of the Nucleus
Figure 26.14 shows the evolution in surface roughness
during film growth, as obtained by spectroscopic ellipsometry (SE), for three hydrogen dilution ratios R of 0,
10 and 20 [26.26]. As is seen in the figure, after the formation of an island, the enforced coalescence of islands
takes place, which results in a smooth surface under
50
Roughness (Å)
R = 20
40
30
R=0
20
10
R = 10
0
100
101
102
103
104
Thickness d (Å)
Fig. 26.14 Evolution in surface roughness, as measured us-
ing spectroscopic ellipsometry, during film growth for three
different Rs
587
H
Si
Dangling bond
Part C 26.2
growth occurs layer-by-layer via an alternating sequence
of thin amorphous film growth and hydrogen plasma
treatment. Several monolayers of amorphous silicon are
deposited, and these layers are exposed to hydrogen
atoms produced in the hydrogen plasma. This procedure
is repeated several tens of times in order to fabricate the
proper thickness to be able to evaluate the film structure. The absence of any significant reduction in film
thickness during the hydrogen plasma treatment is difficult to explain using the etching model, and so the
chemical annealing model was proposed, as schematically shown in Fig. 26.13. During the hydrogen plasma
treatment, many atomic hydrogens permeate through the
subsurface (the growth zone), giving rise to the crystallization of an amorphous network through the formation
of a flexible network without any significant removal of
Si atoms [26.23, 24].
These three models have been carefully examined,
and the merits and drawbacks of each model have been
discussed [26.24, 25].
More microscopic observations have recently been
reported, based on the use of in situ diagnostic techniques, and a detailed mechanism for the formation
process of µc-Si:H has been proposed.
26.2 Film Growth on a Surface
Absorbance difference (arb. units)
0
– 0.001
Si-H2(Sid)
complex
– 0.002
– 0.003
1800
Si-Hx bulk
1900
2000
Si-Hx surface
2100
2200
2300
Wavenumber (cm–1)
Fig. 26.15 Surface infrared absorption spectrum from the
film just before nucleus formation, showing the appearance
of the Si−H2 −d complex whose structure is also shown
µc-Si:H-growth conditions (R = 20). After the smooth
surface is obtained (formation of the nucleus is confirmed at this point in time), surface roughness is then
enhanced due to an orientation-dependent crystal growth
rate [26.27]. As soon as the smooth surface appears, a
particular surface absorption band is observed in the infrared absorption spectrum, as measured using the in
situ attenuated total reflection technique (ATR) during
film growth [26.27, 28]. Figure 26.15 shows the surface
infrared absorption spectrum, showing the presence of
specific bands at 1897 cm−1 and 1937 cm−1 together
with the usual Si−Hx surface and bulk absorption bands
(which occur between 2000 cm−1 and 2150 cm−1 ). This
new absorption band is assigned to the SiH2 (Sid) complex, which is also sketched in Fig. 26.15. Note that the
number density (absorption intensity) of the SiH2 (Sid)
complex is found to be proportional to the magnitude
of the internal stress in the film just before these complexes appear. A nucleation model has been proposed
based on the experimental facts mentioned above. The
enforced island coalescence due to the enhanced surface diffusion of SiH3 gives rise to an internal stress
involving many strained Si−Si bonds in the amorphous
588
Part C
Materials for Electronics
Part C 26.2
incubation layer [26.27]. Atomic hydrogen attacks the
strained Si−Si bonds, forming SiH2 (Sid) complexes
on the film-growing surface. These complexes provide
structural flexibility, which enables structural order to
be obtained via successive Si−SiH3 bond formation at
this site; in other words it acts as a prenucleation site on
the film-growing surface.
Epitaxial-Like Crystal Growth
Figure 26.16 shows a cross-sectional transmission electron microscope (TEM) image of typical µc-Si:H films
deposited on a glass substrate [26.29]. In the figure,
epitaxial-like crystal growth is clearly observed to occur
from the nucleus. It is a well-known fact that epitaxial
crystal growth occurs only when the surface diffusion
length of the film precursor is sufficiently long. In order
to investigate the origin of the enhanced surface diffusion of SiH3 during the formation of µc-Si:H, an isotope
labeling experiment has been carried out [26.30]. D2
was used as the source gas material (D2 SiH4 ) instead
of H2 SiH4 during film growth under constant substrate
temperature conditions, and the number densities of both
D and H incorporated into the resulting film are measured by infrared absorption spectroscopy in order to
estimate the degree of H-to-D exchange reactions, which
mostly occur on the film-growing surface.
Figure 26.17 shows the Raman crystallinity Ic /Ia
plotted against the number density ratio of the D/H
incorporated in the resulting films. As is clearly from
the figure, µc-Si:H is obtained only when D/H exceeds
a critical value, indicating that the D/H exchange reaction is required to some extent for the formation
of µc-Si:H. The D/H exchange reaction involves two
steps:
1. Atomic D reaching the film-growing surface abstracts bonded H on the surface (the Eley-Rideal
reaction), forming HD and providing a Si dangling
bond site;
2. A recombination reaction occurs between the dangling bond and another atomic D adsorbed onto the
surface, forming a Si−D bond.
These abstraction and recombination reactions are
known to be strong exothermal reactions, with 1.4 eV
and 3.1 eV of energy released, respectively. Considering
the highly exothermal nature of the D/H exchange reaction, and the existence of a threshold value for the degree
of D/H exchange reactions required for the formation
500 Å
Fig. 26.16 Cross-sectional transmission electron microscope image of typical µc-Si:H deposited on a glass
substrate
12
Ts = 250 °C
Raman crystallinity (Ic / Ia)
a – Si: H
µ c – Si: H
10
8
6
4
2
0
0.1
1
10
D/H
Fig. 26.17 Raman crystallinity plotted against the number
density ratio of D/H incorporated into the resulting films
prepared from D2 SiH4 plasmas
of µc-Si:H, as shown in Fig. 26.17, “local heating” on
the film-growing surface is believed to play a major
role in the enhanced surface diffusion of the film precursor (SiH3 ) during both nucleus formation (strained
bond formation and SiH2 (Sid)-complex formation) and
the epitaxial-like crystal growth associated with µc-Si:H
film growth [26.17, 24, 28, 31].
Amorphous and Microcrystalline Silicon
26.3 Defect Density Determination for a-Si:H and µc-Si:H
589
26.3 Defect Density Determination for a-Si:H and µc-Si:H
26.3.1 Dangling Bond Defects
Figure 26.18 shows the dangling bond defect density
in the resulting a-Si:H as a function of substrate temperature. This dependency of the dangling bond density
on the substrate temperature has been explained by taking into account the steady-state dangling-bond density
on the film-growing surface. The structural properties
of the resulting film are generally largely determined
by the steady-state surface properties during the thin
film growth process, because the surface formed at any
given instant is incorporated into the bulk in the next instant due to the successive layering nature of the film
growth [26.32, 33]. The steady state number density of
surface dangling bonds is determined by the balance between the rate of generation of dangling bonds and their
annihilation rate.
At low substrate temperatures, surface dangling
bonds are produced by abstraction reactions of surface H atoms with SiH3 ; the reaction rate for this is
almost independent of substrate temperature, since the
reaction rate is affected by both the residence time of
1019
Defect density (cm3)
SiH3 at the H-covered site (longer residence times occur
at lower temperatures) and by the abstraction reaction
rate (slower rates occur at lower temperatures). The surface dangling bond is saturated by SiH3 diffusing across
the surface (this is dangling bond annihilation), which
exhibits a slower rate at low temperatures due to the
slower surface diffusion of SiH3 , which results in high
dangling bond density on the steady-state film-growing
surface. As a consequence, the dangling-bond density
in a-Si:H grown at low substrate temperatures can be as
high as 1019 cm−3 . When the substrate temperature is
increased during film growth, the steady state number
density of surface dangling bonds is reduced drastically
due to the combination of the temperature-independent
H-abstraction reaction and the thermally enhanced surface diffusion of SiH3 , which gives rise to a minimum
defect density of 1015 cm−3 in the resulting a-Si:H at
a substrate temperature of ≈ 250 ◦ C. The increased dangling bond density of the a-Si:H prepared at substrate
temperatures higher than 350 ◦ C is explained by the increased rate of generation of steady-state dangling bonds
due to the addition of a new term in the generation rate of
these bonds associated with the thermal removal of surface H, although its annihilation rate strictly increases
with increasing substrate temperature [26.33, 34].
We note here that a remarkable increase in the contribution from short- lifetime species such as SiH2 , SiH and
Si, which show high reactivity and no diffusion on the
film-growing surface, is observed upon the depletion of
SiH4 -parent molecules, which causes an increase in the
1017
1018
Defect density (cm–3)
1016
1017
1015
250°C
400°C
450°C
16
10
1015
0
100
200
300
400
500
600
Substrate temperature (°C)
Fig. 26.18 Dangling bond defect density in the resulting aSi:H films as a function of the substrate temperature during
film growth
1014
0.1
1
10
Deposition rate (A / s)
Fig. 26.19 Dangling bond defect density in a-Si:H films
plotted against their deposition rates for three different
substrate temperatures
Part C 26.3
One of the most important structural properties of a-Si:H
and µc-Si:H for device applications is their dangling
bond defect densities, because each dangling bond creates a localized deep state in the band gap of the material,
which acts as a recombination center for photoexcited
electrons and holes, although the free carrier mobility is
another important property in semiconductors.
590
Part C
Materials for Electronics
Part C 26.4
number density of dangling bond defects in the resulting films through an enhancement of the dangling bond
generation rate and a reduction of the dangling bond
annihilation rate on the film-growing surface [26.20,32].
Based on our understanding of the defect density
determination reaction during film growth, several trials
have attempted to control the defect density in a-Si:H.
The steady state defect density on the film-growing surface could be reduced when the growth rate was much
faster than the thermal H removal rate in the substrate
temperature range above 350 ◦ C, where the steady state
defect density is mainly determined by the thermal H
removal process. Figure 26.19 shows the number density of dangling bond defects in a-Si:H films plotted
against their growth rate. As expected, the defect density in a-Si:H shows no growth- rate dependence when
the film is prepared at substrate temperatures lower than
300 ◦ C, whereas the defect density monotonically decreases with increasing growth rate when the film is
deposited at 400 ◦ C and 450 ◦ C, and a defect density of
1014 cm−3 has been demonstrated [26.35].
26.3.2 Dangling Bond Defect Density
in µc-Si:H
Figure 26.20 shows typical dangling bond defect densities in µc-Si:H as a function of substrate temperature
together with those in its a-Si:H counterpart [26.33, 35,
36]. As seen from the figure, the substrate temperaturedependent dangling bond defect density in µc-Si:H in the
temperature range above 300 ◦ C shows a similar trend to
that for a-Si:H, indicating that the defect determination
reaction on the film-growing surface is identical for both
µc-Si:H and a-Si:H growth in this temperature range; in
other words, the increase in defect density with increasing substrate temperature is controlled by the thermal
1019
Defect density (cm–3)
a – Si: H
1018
1017
1016
µ c – Si: H
1015
0
100
200
300
400
500
600
Substrate temperature (°C)
Fig. 26.20 Dangling bond defect density in µc-Si:H films
as a function of substrate temperature during film growth
in comparison to that in a-Si:H
H removal process on the film-growing surface. On the
other hand, the number density of defects in µc-Si:H prepared at low substrate temperatures shows much lower
values than those in a-Si:H. Considering that the defect density in the resulting film is largely determined
by the defect density on the steady-state film-growing
surface, the much lower defect density in µc-Si:H is
caused by an increase in the defect annihilation rate on
the film-growing surface due to the local heating from
H-exchange reactions that occur during the course of
µc-Si:H film growth.
It should be noted here that the dangling bond defect
density in the resulting µc-Si:H is strongly influenced by
the ion bombardment during film growth and by the collisions of crystals growing from different nuclei, which
are different aspects to those that are important during
a-Si:H growth.
26.4 Device Applications
a-Si:H and µc-Si:H are highly promising materials applicable to electronic or optoelectronic thin-film devices
such as thin-film transistors (TFT), position sensors,
color sensors, solar cells, etc. [26.1, 2].
A thin-film transistor array with a-Si:H active layer
has been developed for switching devices in liquid crystal displays (LCD), and large-area LCDs more than 40 in
across have already been made commercially available
based on this technology. A laser crystallization technique has been developed in order to increase the carrier
mobility of a-Si:H-based TFTs and thus reduce the ac-
tive area of the transistors in the LCD. LCDs with a-Si:H
or laser-crystallized thin-film Si-based transistor arrays
are widely used as flat panel displays in televisions and
monitors, where they are in competition with plasma display panels (PDP) and other flat panel display systems.
If the carrier mobility in as-deposited µc-Si:H is drastically improved through the enhanced control of the film
growth process, µc-Si:H will be widely used not only for
thin-film transistors but also for signal-scanning devices
such as charge-coupled devices monolithically arranged
in the periphery of the LCD.
Amorphous and Microcrystalline Silicon
26.5 Recent Progress in Material Issues Related to Thin-Film Silicon Solar Cells
any photoinduced degradation, although the conversion
efficiency naturally deteriorates due to the reduced absorption of sunlight. The fabrication of a tandem-type
stacked solar cell structure consisting of a top cell
with a thin a-Si:H layer and a bottom cell containing
narrow-gap materials such as a-SiGe : H, µc-Si:H and
µc-SiGe:H has been proposed as a promising way to
overcome photoinduced degradation and so to achieve
high conversion efficiency in thin-film silicon-based solar cells. Initially, a-SiGe:H was adopted for the bottom
cell material [26.38]; however, this material also shows
severe photoinduced degradation. Recently, µc-Si:H or
µc-SiGe:H have been proposed as promising candidates
for bottom cell materials, because these materials do
not exhibit photoinduced degradation [26.39,40]. In this
proposal, high rates of growth of those materials are
crucial to the low-cost fabrication of tandem-type solar
cells, since µc-Si:H and µc-SiGe:H undergo largely indirect optical transitions. Therefore, urgent material issues
for the realization of low-cost/high-efficiency thin-film
silicon-based solar cells include the need to improve the
photoinduced stability of high-quality a-Si:H and the
need to achieve a high rate of growth of device-grade
µc-Si:H or µc-SiGe:H.
26.5 Recent Progress in Material Issues
Related to Thin-Film Silicon Solar Cells
26.5.1 Controlling Photoinduced
Degradation in a-Si:H
A relationship between the degree of photoinduced
degradation and the dihydride bonding (Si−H2 ) density
has been reported in a-Si:H prepared under a variety of
deposition conditions, where the substrate temperature,
plasma-excitation frequency, gas-flow rate, hydrogen dilution ratio, working pressure, power density, etc. have
all been varied [26.41]. Figure 26.21 shows the degree of
photoinduced degradation, defined as the difference in
the fill factors of photo-I–V characteristics of Ni-a-Si:H
Schottky diode before and after light soaking plotted
Fig. 26.21 Relationship between the degree of photoinduced degradation and the dihydride (Si−H2 ) content in
a-Si:H prepared under various deposition conditions. The
star symbol represents for a-Si:H prepared under conditions
of a reduced ratio of contributions from higher silanerelated species versus the SiH3 contribution at a substrate
temperature of 250 ◦ C
against the Si−H2 density in a-Si:H film, as measured
by infrared absorption spectroscopy. Furthermore, it has
been suggested from mass spectrometric results that the
Si−H2 density in the resulting a-Si:H is strongly in-
0.20
Degree of photo-induced
degradation FFinitial – FFafter
0.15
0.10
0.05
0
0
2
4
6
8
10
Dihydride content CH (Si – H2)(at. %)
Part C 26.5
Thin-film Si-based solar cells has also been widely
expected to provide low-cost photovoltaics. Actually, aSi:H-based solar cells have already been widely used
in pocket calculators, and now large-area solar cells
are being developed for electricity generation. One big
advantage of using a-Si:H for solar cell applications
is its large optical absorption coefficient compared to
single-crystalline or polycrystalline silicon counterparts,
resulting in indirect optical transition properties, and
so a thickness of less than 1 µm is enough to absorb
sufficient sunlight for electricity generation when using a-Si:H-based solar cells. Low-temperature processes
using PECVD are also advantageous in terms of reducing the cost of producing a-Si:H-based solar cells.
However, there is a well-known phenomenon that occurs in a-Si:H called photo-induced degradation [26.37],
where the conversion efficiency of a-Si:H-based solar cells, usually ≈ 10%, is degraded to less than 8%
after prolonged exposure to light. Recombinations photogenerated electrons and holes are believed to trigger
this photoinduced degradation in a-Si:H; therefore, increasing the field in the a-Si:H-based solar cell has
proved an effective way to reduce the degradation. In
fact, 0.1 µm-thick a-Si:H-based solar cells do not show
591
592
Part C
Materials for Electronics
Si4H9
SiH3
Part C 26.5
Fig. 26.22 Surface reaction image of the incorporation of
Si−H2 bonds into the resulting a-Si:H due to the contributions from higher silane-related species such as Si4 H9
20
Current density (mA / cm2)
15
10
2 nm / s
Area: 0.25 cm2
Jsc: 16.1 mA / cm2
Voc: 0.854 V
FF: 0.595
␩ : 8.17 %
5
0
0
0.2
0.4
0.6
0.8
1.0
Voltage (V)
Fig. 26.23 Photo-I–V characteristics of a stable n–i–p aSi:H-based solar cell fabricated at a high growth rate of
2 nm/s
creased by the contributions from higher silane-related
species (HSRS) such as Si4 H9 when the substrate temperature is kept constant. The contribution of these
HSRS to the film precursor (SiH3 ) has been theoretically analyzed using a couple of gas-phase reaction-rate
equations [26.42]. This analysis predicted that the contribution ratio is a complex function of the electron
temperature in the plasma, the electron density in the
plasma, the gas temperature, the hydrogen dilution ratio
R, and the gas residence time during film growth. Based
on our understanding of the network structure (Si−H2
bonding configuration) responsible for the photoinduced
degradation in a-Si:H and the chemical species (HSRS)
responsible during film growth (shown schematically
in Fig. 26.22), a guiding principle for obtaining highly
stabilized a-Si:H has been proposed [26.42, 43].
By following the guiding principle, a-Si:H with minimized Si−H2 density in the network has been prepared
by adjusting the plasma parameters during film growth
under high growth rate conditions, and an a-Si:H-based
solar cell showing a stable conversion efficiency of 8.2%
(Fig. 26.23) has been fabricated at a high growth rate of
2 nm/s [26.41].
Moreover, a-Si:H containing a Si−H2 density of 0%
has been successfully prepared using a triode reactor at
substrate temperatures as low as 250 ◦ C by making use
of the difference in the gas phase diffusion coefficients
of SiH3 (light) and HSRS (heavy) during film growth.
26.5.2 High Growth Rates
of Device-Grade µc-Si:H
In the past, conventional high hydrogen dilution methods performed at relatively low working gas pressures
(several tens of mTorr) have been used to obtain devicegrade µc-Si:H (with low dangling bond defect densities
of ≈ 1016 /cm3 ) [26.17]. Recently, a simple concept for
preparing device-grade µc-Si:H at a high growth rate
has been proposed, known as the high-pressure depletion (HPD) method [26.21, 44]. The production rate of
the film precursor SiH3 in the plasma, which is proportional to the growth rate of µc-Si:H, is determined
by the product of the electron density and the number density of SiH4 . To increase the production rate of
SiH3 in the plasma, a high power density (radio frequency, RF, or very high frequency, VHF) and high
partial pressures are needed, because the electron density is basically a function of the power density applied
1018
Defect density Ns(cm–3)
(A) Conventional
(B) High pressure depletion
1017
1016
1015
0
10
20
30
40
50
60 70 80 90
Deposition rate (A / s)
Fig. 26.24 Relationships between the dangling bond de-
fect density in µc-Si:H films and their deposition rates
for (A) µc-Si:H prepared using the conventional lowpressure regime and for (B) µc-Si:H prepared using the
high-pressure depletion (HPD) method. The star symbol
represents µc-Si:H prepared under HPD conditions with
the novel cathode design
Amorphous and Microcrystalline Silicon
26.5 Recent Progress in Material Issues Related to Thin-Film Silicon Solar Cells
25
Current density (mA / cm2)
20
15
2 nm / s
Area: 0.25 cm2
Jsc: 22.3 mA/ cm2
Voc: 0.52 V
FF: 0.70
␩ : 8.1%
10
5
0
Part C 26.5
to the plasma and the number density of SiH4 is proportional to the partial pressure. However, hydrogen atoms
(the chemical species responsible during the formation
of µc-Si:H) are strongly scavenged by SiH4 molecules
during their transportation from their production site to
the film-growing surface. In order to enhance the survival of the hydrogen atoms, it is has been suggested
that the SiH4 molecules should be depleted by applying high power density to the plasma. As conditions of
high total pressure are also useful for decreasing the effects of ion bombardment during film growth through
the reduction of the electron temperature in the plasma,
high working pressures along with SiH4 depletion conditions (HPD) have recently been popularly adopted for
the high-rate growth of device-grade µc-Si:H.
Figure 26.24 shows the number density of dangling
bond defects, as measured by electron spin resonance
(ESR), in the resulting µc-Si:H as a function of growth
rate [26.45]. When the conventional low pressure regime
is used for the growth of µc-Si:H, the defect density increases exponentially with increasing growth rate, as
seen in the figure (see A). This is caused by both an increase in the ion bombardment (due to the high power
density) and an increase in the contributions from shortlifetime chemical species such as SiH2 , SiH and Si due
to the reduced ability of those species to react with
SiH4 and H2 (due to the low pressure and SiH4 depletion) during film growth. On the other hand, the slope
in Fig. 26.24 becomes shallower when HPD conditions
are used during the growth of µc-Si:H, as shown in the
figure (see B), which illustrates the usefulness of the
HPD method for obtaining high-quality µc-Si:H at high
growth rates.
The validity of the HPD method has also been
demonstrated during the fabrication of µc-Si:H-based
solar cells. Figure 26.25 shows the photo-I–V characteristics of a µc-Si:H-based p–i–n single-junction solar cell
prepared at a high growth rate of 2 nm/s using the HPD
method, which exhibits a reasonably high conversion
efficiency of 8.1% [26.46].
However, HPD conditions require that the spacing
between the cathode and anode is reduced in the conventional capacitively coupled plasma reactor, giving
rise to nonuniform plasma production and nonuniform
film growth. To overcome this problem encountered
when using HPD conditions, the structure of the cathode
surface has been designed to produce uniform plasma
593
0
0.1
0.2
0.3
0.4
0.5
0.6
Voltage (V)
Fig. 26.25 Photo-I–V characteristics of a p–i–n µc-Si:Hbased solar cell fabricated at reasonably high growth rate
of 2 nm/s under HPD conditions
1 mm
2
1
Fig. 26.26 Photograph of the surface structure on the cathode de-
signed for the production of high-density/uniform plasmas
production even in a large-area parallel plate electrode
configuration. Figure 26.26 shows the structure of this
novel design of cathode. A multitude of holes (hollows)
with interconnecting slots are arranged on the cathode
surface, which cause strong coupling between the highdensity plasmas produced in each hole where source gas
injection is performed. Using this newl type of cathode, quite high growth rates (more than 8 nm/s) have
been obtained along with reasonably low defect densities in the resulting µc-Si:H, as shown by star symbol
in Fig. 26.24 [26.47].
594
Part C
Materials for Electronics
26.6 Summary
Part C 26
In this chapter, the processes involved in the growth of
a-Si:H and µc-Si:H from SiH4 and SiH4 H2 plasma have
been interpreted in detail. The defect density determination reaction that takes place on the film-growing surface
was discussed in order to obtain clues that may lead to
enhanced optoelectronic properties in those materials.
The recent status of work done in the fields of thin-film
transistors and solar cells was reviewed, as these are
the main device applications of those materials. Recent
progress in resolving material issues related to solar cell
applications were also described.
Finally, we note here that the concepts used in and
our understanding of the film growth process mentioned
here are widely applicable to other processes, especially processes where thin films are grown from reactive
plasmas.
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