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06 chapter 2

6&4P78& TWO
11.1: Historical Background
For over two decades or so, hydrogenated amorphous silicon germanium
(a-SiGe:H) alloy films have been of much interest as a promising material for the
production of solar cells. Great efforts have been made to improve the optoelectronic
properties of a-SiGe:H alloy films, desired as a small energy band gap component of
multi-junction solar cells and other devices employing amorphous hydrogenated
semiconductors[l-6]. Also a-SiGe:H is widely used for the production of multijunction solar cell, since the optical energy band gap (Eopt) can be easily tuned to
match the wide range of the solar radiation spectrum. It has been reported that with
increasing germanium content (CGe) the Eopt of these alloys shifts to lower energies.
But the relationship between each property of the alloy film has not yet been well
understood compared with relatively detailed comprehension that exists for
hydrogenated amorphous silicon (a-Si:H)[7,8]. This was mainly due to the complexity
caused by the addition of another component germanium (Ge) to a-Si:H. The actual
composition of a-SiGe:H alloy film has been found to depend strongly on the detailed
deposition conditions. Therefore, optoelectronic and structural properties of a-SiGe:H
alloy films deposited under various conditions have been studied using quite number
of investigations. Most of these studies have dealt with hydrogenated samples
deposited by glow discharge technique[2,9-16]. Besides the enhanced optical
absorption for longer wavelengths, the electronic properties which are relevant for
photovoltaic applications are strongly deteriorated by alloying Ge to a-Si:H. These
properties are like Urbach energy, defect density, and microstructural factor increase
whereas the ambipolar diffusion length decreases with a reduced £^,[17-21]. In spite
of the reduction in Eopt, a-SiGe:H alloy film based diodes have been successively
applied as a bottom cells in stacked solar cells[22]. The electronic quality of a-SiGe:H
alloy films has been improved by optimizing deposition conditions in both diode- or
triode- glow discharge reactors[23-25]. Therefore, in a conventional radio frequency
(RF) glow discharge technique, deposition of a-SiGe:H alloy films with different CGe
and narrow Eopt were carried out using silane (SiKU) and germane (GeKU) as a source
gas mixtures. But the results of these films indicate poor optoelectronic properties
compared to those of the best a-Si.H films deposited from glow discharge[26]. The
deterioration of the optoelectronic properties of a-SiGe:H alloy films was attributed to
the preferential attachment of hydrogen (H2) atoms to silicon (Si) rather than Ge[27],
which in turn leads to the formation of a considerable number of dangling bonds on
the Ge atoms. In addition, an increase of the dihydride (Si-H2) groups has been
observed and negatively affects the material properties via the increased structural
disorder[28]. Several attempts have been made to overcome this problem by using
different source gas mixtures either by diluting SiFU and GeFL* with hydrogen (H2) or
by an introduction of fluorine (F) in addition to H2 as a bond terminator[29-33]. Some
improvement in the material quality has been reported in the literature of the
fluorinated material [34]. Figure (1) shows the schematic diagram of a typical mixture
of the source gases utilized in the plasma deposition of a-SiGe:H alloy films. These
gases are diluted either with hydrogen or hydrogen and fluorine to deposit
hydrogenated or hydrogenated and fluorinated SiGe alloy films.
In 1983, Nozawa et al.[29] have discovered that deposition of a-SiGe.H alloy
films by glow discharge from gas mixtures of fluorine and hydrogen (SiF4 + GeF4 +
H2) improved the photoconductivity {aPh) of alloys of EopX « 1.5 eV by about an order
of magnitude. This was with a speculation that F might be a suitable bond terminator
in the alloys than in unalloyed Si[35]. Thereafter, Mackenzie et al.[3] have established
the existence of significant differences in a-SiGe:H alloy films microstructure when
they are made from a source gas mixture of SiF4 and GeF4 diluted with hydrogen
rather from a mixture of SiH4 and GeH4. However, deterioration in optoelectronic
properties of Ge-rich alloys of a-SiGe:H has been attributed to the changes in the
band-structure and an increased nonuniformity of the film microstructure[3]. In
general development of a-SiGe:H alloy material revealed a common feature which
shows sharp deterioration in the photovoltaic quality of the material of Eopt < 1.5 eV.
Because of that several attempts have been made to improve the material quality using
various techniques. Zeman et al.[36] have found that hydrogen dilution technique was
of major importance to grow structurally dense material with highly improved
optoelectronic properties suitable for device applications. The main objective of the
above techniques is to energize the heavy and sticky Ge radicals in order to reduce
1 -Schematic
of Si and
diagram of
Ge volatile
SiGe alloy
source gas
Although a lot of progress has been made as far as the important technological
aspects of a-SiGe:H alloys are concerned, many basic properties of these materials are
still not well understood. This is not surprising if one considers the large sample space
spanned by the different compositions, preparation methods, deposition conditions,
and doping levels. Several attempts have been made in the literature to explain the
degradation of the film quality with the deposition conditions. The interpretations are
based on the existence of various structural and compositional inhomogeneities[3739]. The structural inhomogeneity was represented by the heterogeneous network
while compositional inhomogeneity lead to two types of film microstructure
associated with clustered hydrogen and Ge concentration fluctuations. Matsuda[40]
have introduced that the network structure of a film is principally determined by its
surface growing process and low density films are likely to grow if the surface
diffusion coefficients of the adsorbed radicals are too small to find out their
energetically(fkvorabie^sites. According to this argument the structure of a-SiGe:H
alloys can be discussed in terms of the surface diffusion coefficient of adsorbed
Various studies on the optoelectronic, structural, and spin properties of
a-SiGe:H alloys deposited by a glow discharge of SiHt, GeH*, and H2 gas mixtures
were carried out using different characterization techniques over a wide range of
deposition conditions. The most commonly used techniques to evaluate the material
properties are the measurements of dark conductivity (a</), photoconductivity (aph) and
subband gap absorption. Ray et al.[41] have observed the increase in oph of the plasma
enhanced chemical vapor deposited (PE-CVD) a-SiGe:H alloy films when they are
deposited at slow rates by controlling the gas flow rates of SLH4 and GeTLt, and
chamber ^pressure. Also they have reported that for Ge-rich films with Eopt < 1.5 eV
low deposition rate and hydrogen dilution of the gas mixture are necessary to improve
the material quality. Using the super chamber method, Haku et al.[42] have obtained a
high photosensitivity {aph^o^) of 106 for an a-SiGe:H alloy films containing low
impurity concentration. Recently Xu et al.[43] have succeeded to fabricate high quality
narrow band gap a-Sio.5sGeo.42:H alloy films by controlling the bonded hydrogen
content and its configurations using the nanometer deposition or hydrogen plasma
annealing method.
The subband gap absorption and other optical properties of an a-SiGe:H alloy
films have been measured using various techniques like photothermal deflection
spectroscopy (PDS), constant photocurrent method (CPM), luminescence, etc. These
measurements give information about the Urbach energy and the near midgap deep
states. By using PDS, Stutzmann et al.[5] have observed a strong increase in the
absorption with increasing Ge content, this increase was accompanied with a
broadening of the exponential Urbach tail due to the presence of defects in the
mobility gap. In addition to this the Eopt decreases from a-Si:H to a-Ge:H. Following
Eopt a parallel shift in the band-to-band luminescence energy accompanied with a
decrease in luminescence intensity when going from a-Si:H to a-Sio.4Geo.6:H alloys
have been observed due to the thermalization losses[5]. Figure (2) shows the optical
The hydrogen content (CH) and its bonding configuration in a-SiGe:H alloy
films were studied using both nuclear magnetic resonance (NMR) and fourier
transform infrared spectroscopy (FTIR). With these techniques strong correlation has
been found between the total hydrogen concentration and the alloy composition. The
total hydrogen concentration in the alloy material was predominantly determined by
substrate temperature (7;). NMR measurements shows that about 10 at.% hydrogen
with compositions close to 50% Si and 50% Ge are accompanied with presence of
maximal chemical disorder in the material, and all of the additional hydrogen was
incorporated in the clustered phase (broad proton resonance)[5]. While the FTIR
spectra exhibits Ge-H, Si-H, and Si-Hfe as the main hydrogen bonding configurations.
Figure (3) shows data for variation of CH with Ts. Disregarding the existence of
atomic vacancies or voids and the transition of the film rigidness caused by the
difference in C«[45-48], the widening of the Eopt with CH is due to the valence band
energy being pushed down because the Si-H (3.35 eV) or Ge-H (3.0 eV) bonds are
stronger than Si-Si (1.83 eV) or Ge-Ge (1.63 eV) bonds. Therefore, switching Si-Si
Fig . 2 : Optical absorption coefficient spectra
for a - S i G e - H alloy films deposited with
different alloy compositions-
bonds to Si-H bonds or Ge-Ge bonds to Ge-H bonds[49] caused an increase in bond
energy per hydrogen atom.
The presence of dangling bonds as the main localized defects in the mobility
gap of the a-SiGe:H alloy films was estimated from electron spin resonance (ESR)
measurements. The degradation of the optoelectronic properties of a-SiGe:H alloy
films was attributed to the presence of unwanted preferential formation of Ge-Like
dangling bonds closely linked to the preferential incorporation of Ge atoms from the
gas discharge. Regarding this, Stutzmann et al.[5] have observed that the Ge-like
dangling bonds appear for Ge concentration of « 30% and in the Ge-rich alloys only
Ge dangling bonds are seen with a density of « 5xl017^cm*3j As far as the electronic
structure of the valence-band tail in a-SixGei.x:H, is concerned, the deepest valenceband tail states are preferentially caused by localized bonding orbitals of weak Ge-Ge
bonds, whereas Si-Ge and Si-Si bonds lead to more shallow states[5].
The morphology of a-SiGe:H alloy films has also been studied by various
groups using transmission electron microscopy (TEM), scanning electron microscopy
(SEM), and scanning tunneling microscopy (STM). Using TEM Mackenzie et al.[3]
have observed for a-Sio.sGeo.s'.H deposited at Ts = 230°C a distinct two-phase
microstructure (island tissue structure), as shown in Figure (4), consisting of
noncoalescing island regions of high electron density surrounded by lower density
tissue material. But the photoresponse of the material deteriorates with an increase in
the structural inhomogeneity when Ts is decreased below the optimum value[2]. Also
the morphology of etched a-SiGe:H alloy films was investigated by Zeman et al.[36]
using SEM micrographs, they observed that alloying of a-Si:H with Ge leads to an
increased microstructure in the films. In addition to the morphological studies, Haku et
al.[42] have used STM to study the roughness and film rigidity of intentionally
contaminated a-SiGe.H alloys. Their observations show deterioration of the film
rigidity with increased impurity concentration. Of course, high impurity concentration
causes a decrease in the photosensitivity (ap>/<Jd) due to the increase in defect density.
A lot of progress has been made to study the microstrucrural properties of
a-SiGe:H alloy films. Using small angle X-ray scattering (SAXS), Jones et al.[50]
have observed a sharp change in the anisotropic character of the SAXS for a-SiGe:H
T s (°C)
Hydrogen content ( C H ) as a function
(T s ) f o r a-SiGe:H alloy
Fig . U '• D i s t i t i n c t
alloys with CGe < 0.40. This was in consistent with a transition to a columnar-like
microstructure above CGe - 0.2. Therefore strong evidence was obtained for a direct
link between the degraded optoelectronic properties and the increased heterogeneity
associated with the microstructural transition.
From the previously reported studies, the apparent consensus regarding the
best a-Sii.xGex:H material may be summarized as follows: The continuous random
network (CRN) structure of a-Sii.xGex:H alloys is essentially the same as that of
a-Si:H, the morphology on a scale of 1-10 nm is very similar; the H- incorporation and
diffusion demonstrate the same properties; the defect structure of voids, islands, tissue
and columns are similar to that of a-Si:H and they may differ only in the quantative
details; and the electronic band structure, valence and conduction band densities of
states as a function of energy, the tails to these bands of more localized states, the
defect levels such as those associated with dangling bonds differs only in their scale.
Thus the structural and optoelectronic properties of the best a-Sii-xGx:H alloys should
be comparable to a-Si:H for those parameters judged of high quality. Nevertheless, the
optoelectronic response is consistently poorer, clearly less conducive to use in devices,
and without a clear identification as to whether the responsible entity is the basic
transport mode in a homogeneous low point-defect density structure, or a high pointdefect density , and /or inhomogeneous morphology causing lower carrier mobility
and enhanced recombination.
It has been discovered that thin amorphous silicon based alloy semiconductors
can be prepared
in multilayered
with crystalline
properties[51]. These multilayered structures are deposited from alternative layers of
different alloy materials and they have been the focus of increasing interest due to
their potential technological applications[51-53] in solar cells and photoreceivers and
other devices. Additionally, the field of amorphous semiconductor multilayers has
developed so rapidly that a variety of interesting phenomena similar to those of
crystalline multilayers can be found. In the literature different
configurations like a-Si:H/a-SiC:H, a-Si:H/a-SiO.H, and a-Si:H/a-Ge:H have been
studied. The optical band gap of the a-Si:H layers in the multilayered structures has
been found as a consequence to increase as its thickness decreases[51,54]. In 1985,
Maley and Lannin[55] have observed the heterostructure effects on the phonons of
measurements. Raman spectra imply increased bond-angle disorder associated with
interfacial induced modifications of network formation with decreasing thickness of
a-Si:H layer. Also they observed that the interfacial bonding constraints in amorphous
superlattices result in local disorder that qualitatively differs in character from that of
crystalline systems. The amorphous silicon based alloy multilayers also have served as
an experimental system in which the interfaces between the layers can be studied[56].
Several characterization techniques have been used to study the multilayered
structures. Such techniques are like photothermal deflection spectroscopy (PDS)[57],
optical absorption[58], Raman scattering[55], and persistent photoconductivity[59-61].
11.2: Hydrogen Chemical Potential
The hydrogen density of states (HDOS) model can describe the kinetics of the
hydrogen incorporation in a-Si:H and its related alloys. In this model, the HDOS in
which the sites in the material where the hydrogen resides can be regarded as
traps[62]. These traps can be distinguished into two types, shallow as well as deep
traps. The shallow traps being the bond-center and interstitial sites in the network
while the deep traps being the dangling bond sites. Since the hydrogen incorporation
takes place during the growth of the film which is governed by the equilibrium
between the hydrogen chemical potentials (/XH) in the film and the gas phase.
Therefore JUH can be determined by using the deep traps energy level Ea, substrate
temperature (7^), density of the occupied deep traps by hydrogen (NH), and the total
deep trap density (A^f) via the following expression;
( N
M„=Edl+ KTS In — ^
During the film growth /J.H is situated between the deep trap energy level (Edi)
and the shallow trap energy level (Est). The JJH of hydrogen incorporated in the film
increases with Ts. The JUH in the film and in the gas should be equal during film
deposition. Hence the traps occupancy must be decreased by an increase in Ts, and the
hydrogen content (C#) in the film be decreased. Accordingly, the high hydrogen
concentration in the films deposited at lower deposition temperature can be attributed
to the higher //# of atomic hydrogen[65].
11.3: Hydrogen Density of States
It is well known that a-Si:H binary alloys are classified on the basis of
hydrogen bonding strength to the alloy element compared to a-Si.H. Alloys like
a-SiC:H, a-SiO:H, and a-S£N:H have stronger hydrogen bonds while a-SiGe:H alloys
have weaker bonding. These hydrogen bonds to the alloy elements have different
binding energies, as can be seenfromTable (1).
Table (1)
Bond strengths of various hydrogen bonds
Energy (eV)
Figure (5) illustrates the hydrogen density of states (HDOS) of a-SiGe.H alloy films
and shows the different hydrogen bonding states. EM is the effective migration energy
of hydrogen in a-SiGe:H and ER is the redistribution energy, which defined by the time
available for structural change. In the HDOS, hydrogen occupying states above ER.
The weaker H bond in a-SiGe:H element implies that the Si-H bonds determine the
structural rearrangements. Earlier reports
shows that
band tails of the
a-SiGe:H alloys are different from a-Si:H[63,64]. The problem arising during the
growth of a-SiGe:H alloys is that the Ge-H bond overlaps the weak bond distribution
and the hydrogen chemical potential /JH lies in a large density of states resulting in a
high defect density. These defects are expected to be Ge dangling bonds because the
Ge-H bond is at higher energy. It is difficult to improve the electronic properties of
a-SiGe:H alloys because of the overlap between the weak bonds and the Ge-H bonds
in the HDOS. The model suggested by Street[65] predicts that the growth of low
defect density material requires a careful control of //#, and that the range of suitable
growth condition is more costrained than for a-Si:H. Hydrogen dilution of the plasma
Density of states
' S i -H^)
5 • Hydrogen
a-SiGe H
of s t a t e s (HDOS)
films showing
and lower temperature deposition may help to raise /JH- Optimum conditions are
expected to be close to those needed to grow microcrystalline material[65].
11.4: Relative Incorporation Probability
For a-SiGe:H alloy films deposited by glow discharge of SiHt and GeHU it is
generally observed by many workers that the CGe in the deposited solid film is
significantly higher than the germanium concentration in the glow discharge
plasma[66-68]. This means that germanium has a higher deposition rate than silicon
under the standard deposition conditions. Therefore the relative incorporation
probability for silicon and germanium in an a-SiGe:H alloy films can be defined as
P(Ge)/P(Si) =[]-
(1/XsoMl - (1/ZvJJ
Where P(Ge) and P(Si) are the incorporation probabilities of Ge and Si respectively
and Xsoi, and zgas refers to the atomic and volume fractions of silicon in the deposited
a-SiGe:H alloy films ( i.e. silane in the deposition plasma ). Large variations in the
composition of the deposited a-SiGe:H alloy films can also occur even for a constant
gas mixtures if the deposition conditions are changed.
11.5: Photconductivity and Photosensitivity
The indicators that characterize the optoelectronic quality of the material are
both the photoconductivity (<jph) and the photosensitivity (aPf/ad). The study of aPh in
amorphous semiconductors allows a closer understanding of the recombination
process of excess carriers. In its simplest form, the aPh can be expressed as
oph = Anen
Where An is the excess photoexcited electron density, fi is the electron mobility, and e
is the electron charge. In the steady state An - Gx, where G is the photoexcitation rate
and r is the electron lifetime, and since G & naF(J - R), where n is the quantum
efficiency, a is the absorption coefficient, F is the photon flux, and R is the reflection
coefficient. The equation (3) can be rewritten as
an = a(J -RJFeijMT
Depending on the recombination kinetics, r is usually a function of the generation rate.
Generally it is found experimentally that
a* *F
where 0.5< y<\.§ is the exponent of the generation rate and depends on the carrier
generation and recombination kinetics. For a-SiGe:H alloy films y=0.9[44].
The photosensitivity is found as the ratio of aPh measured under AMI
illumination (100 mW/cm2) to a.
Universally it was found that the optoelectronic properties of a-Sii-xGex:H
alloy films are much poorer than those of a-Si:H. The quantum efficiency-mobilitylifetime {rjfiv) product is related to the <jph caused by an absorption of G photons/cm3-s
by the equation
= e<
= e/jTjGr
(6 A)
The quantity An(E) is the steady-state extra carrier density at energy E per unit
energy interval and p(E) is the mobility of the carriers at that energy. The integral over
E of An(E) is related to the carrier production rate (the product of G times an
efficiency factor rf) and average recombination time (r) for electrons and holes. The
first bracket in equation ( 6A) above defines effective carrier mobility in the steady
Many researchers have pursued a correlation between the measured oPh and
deduced T]/JT product and the preparation conditions such as substrate temperature,
applied RF power and measured structural and optoelectronic properties[ 1-6,64,6971]. Thus, the low values of photosensitivity is correlated with low preparation
temperature, too rapid deposition rate, etc. Also it is correlated with structural
properties like heterostructure visible in transmission electron microscopy (TEM)[24,37,72], multiple H-bonding resulting in bending modes seen in infrared (IR)
gas evolution at low annealing temperatures[37,72]. Low
photoconductivity is attributed to low density island joining tissue or voids[50]; and
with undesired optoelectronic properties such as high sub-band gap optical
absorption[2-4], high dangling bond density from studies of electron spin resonance
(ESR), low prefactors in the conductivity-inverse temperature relation, miniscule
photoluminescence intensity[2-4], and high gap state densities as inferred from several
types of study[6,73]. Stutzmann et al.[74] suggest^ that the recombination
in a
semiconductor can occur either by bimolecular band-to-band transitions or
monomolecular recombination via defects. It shows that for a large density of
recombination centers (Ns > 1016 cm"3) under illumination with an intensity of 100
mW/cm2, the electron and hole densities are essentially governed by monomolecular
recombination via direct capture at dangling bond defects. Due to direct capture of free
electrons the electron lifetime was estimated as[44],
where Se is the capture cross section for electrons, v, is the thermal velocity, and Nr is
the number of recombination centers and it can be calculated by the Rose model[75].
Kuznestov et alj[44] assumes that Nr is proportional to the dangling bond density. So
they expect that for a-SiGe.H alloy films the aPh should vary proportionally to NS'J.
The electron mobility can be expressed as
where fa is the free electron mobility and /? is the coefficient of proportionality. In
optimized alloys 0 ~ 1, and consequently aph ~ N/ 7 . In alloys grown at low
temperatures 0 < 1 because hydrogen clustering leads to increased defect density and
increased electron scattering in the conduction band. The electron mobility decreases
and changes from the free electron mobility in optimized alloys to the drift mobility in
low temperature a-SiGe:H alloy films. The decrease in mobility can explain why the
experimental results shows a stronger decrease of oph than increase of N, for low
temperature alloys.
11.6: Band Gap
The characterization of amorphous solar cells is seriously affected by the
optical energy band gap (Eopt) of the photovolatiac material. Therefore, choice of the
Eopt for each i-layer in multigap stack cells is essential because the Eopt limits the cell
performance via the Output current and voltage and is also believed to affect the
stability against light exposure. Several reports have studied the influence of film
composition on the Eopt and most of them refer only to the relationship between the
Eopt and the germanium content (CGe)- The dependence of the Eop, on CGe follows the
relationshipf 1,3,5,76-79],
Eopi = aCGe + b (eV)
The coefficient of CGe was estimated to be about -0.78 eV. Figure (6) shows the
dependence of Eopt on CGe[80].
The Eopt of a-SiGe:H alloy films was easily narrowed by changing its
composition. Much effort has been made to obtain high quality a-SiGe:H alloy films.
However, the film quality significantly deteriorates with the narrowing of the
Eopt[24,42,76]. The deterioration of the film qualities in the narrow Eopt region is
caused by the increase of Ge-Ge bonds in the films[81]. Also it has been seen that Eopt
is determined by the combination of Cae and hydrogen content (C#). With this Eopt
seems to have an almost linear relationship with both CGe and CH and it can be
represented by
F i g . 6 • Dependence
of a-SiGe«H
content (CQ2)
of optical
energy band gap
alloy f i l m s
on the germanium
content ( C g e ) • The solid
point [ • ]
in the
E o p t reported
Eopt = 1.3CH - 0.6CGe + 1-4 (eV)
In this equation the absolute coefficient of CG£ was smaller than -0.78 eV[53]. From
Figure (6) part of the change in the Eopt with CGe was ascribed to the variation in C#,
which is known to decrease with an increase in CGe at a fixed substrate temperature
and RF power[82,83]. According to equation (10), results led to high efficiency and
stable multi-junction solar cells[84].
11.7: Persistent Photoconductivity
The persistent photoconductivity (PPC) effect has been widely observed in
III-V materials. PPC is photoinduced conductivity that persists after the termination of
the photoexcitation which has been observed in many semiconductors[85-87]. PPC
has been used to study the metastability of impurities[85], to study the transport
properties in disordered systems[86], and to adjust the two dimensional carrier density
in an interface[88].
Several models have been proposed to explain PPC
phenomena[89-91]. The separation of electrons and holes by the electric fields of
inhomogeneities can cause anomalously long nonequilibrium carrier lifetimes, which
depend in an activated manner on temperature[92]. In this mechanism, a PPC-related
phase transition has been observed, in which the stored charge carriers experience a
phase transition from hopping-conduction to a percolation-transport state at a critical
temperature, and their relaxation behavior was found to follow the stretchedexponential function[86,89]. Queisser and Theodorou[90,93,94] have demonstrated
that the spatial separation of photogenerated electrons and holes by a built-in electric
field from a macroscopic potential barrier due to band bending at surfaces or interfaces
leads to PPC. This model predicts a PPC decay essentially logarithmic in time that is
This mechanism
photoexcitation of electrons from deep level traps, which undergo a large lattice
The decay kinetics of the PPC observed in many disordered materials could be
described by the stretched exponential function, i.e.
bpc(t) = IPPC (0) exp[-(t/rf]
Where 0< /? <1 and IPPc is the persistent photocurrent.
In equation (3) An can be considered as the change in electron concentration
due to PPC by assuming that any change in
would be due to a change in electron
11.8: Conversion Efficiency
Because of the low cost and large area produceability, a-Si:H solar cells have
attracted much attention in the last few years. But there exists two main shortcomings
which obstruct the spread of applications of this kind of solar cells. These are:
(1) The low absorption efficiency in the long wavelength region, which means that a
great part of the solar energy cannot be absorbed and the conversion efficiency is
(2) The instability of the single-junction a-Si:H solar cell, i.e. the conversion
efficiency will be deteriorated under some external reactions such as illumination
and heating.
For further improvement in the conversion efficiency of a-Si solar cells and to
resolve some of these problems, many investigations have been made, out of which the
most meaningful one is the use of multi-junction a-Si:H solar cells[96,97]. Combining
various band gap materials to effectively absorb the different regions of sun light, the
overall efficiency of the solar cell be increased. The generally used narrow band gap
material is a-SiGe:H which earmarked as a promising material for the bottom cell of
high performance multi-junction solar cells. The conversion efficiency of 10.6% for 10
cm x 10 cm integrated-type single-junction a-Si solar cell submodules have been
achieved by combining a number of techniques such as:
1. A high quality i-layer and p/i buffer layer[98],
2. A high quality p-layer of a-SiC[99],
3. A highly textured transparent conducting oxide (TCO) with a high haze ratio and
gentle slope of grains which can yield a uniform p-layer[100].
Haku et al.[42] have obtained a conversion efficiency of 3.34% for a 1 cm
single-junction solar cell using a-SiGe:H intrinsic layer with a low impurity
concentration. This single-junction solar cell have been also used by them as a bottom
cell for a 1 cm2 a-Si/a-Si/a-SiGe triple-junction solar cell and a conversion efficiency
of 11.9% have been achieved under AM-1.5, 100 mW/cm light.
Extensive research has been made during the last fifteen years or so. During
these years a-Si.H technologies have shown considerable progress. This leads to an
improvement in the performance of its solar cells where a conversion efficiency of
12.0% for a single-junction a-Si solar cell of 100 cm2 module has been achieved[101].
The main problems arise with a-Si solar cells is the degradation of efficiency
due to the increasing number of metastable defects in the photovoltaic layer during
light irradiation, or the so-called Staebler-Wronski effect[102]. This effect is probably
caused by the recombination of electron-hole pairs[103]. Therefore multi-junction
solar cells are expected to relieve this problem, because the thickness of each
photovoltaic layer is thinner than for single-junction solar cells. Similarly a rather
strong internal electric field prevents carrier recombination and maintains a higher fill
factor after illumination. In addition to the cell structure, the stability of the materials
themselves must be improved. Taking all these factors into consideration, the
stabilized conversion efficiencies of 3.3% and 10.6% were achieved for an a-SiGe
single-junction and for a-Si/a-SiGe double-junction solar cells respectively[104].
Since the defect density of a-SiGe:H depends on the band gap, the design of
the band gap profile has a strong influence on the spatial defect state distribution and
thus on the optoelectronic properties of the cell. With band gap profiling significant
improvements in the solar cell performance were achieved by a band gap graded
i-layer[105-107]. Guha et al.[108] have fabricated a multi-junction solar cell with an
efficiency of 13.7% using band gap graded a-SiGe:H bottom cell. Recently an
improved efficiency of 14.6% were achieved for a-SiGe.H multi-junction solar
cells[109]. This efficiency is the highest value reported to date for an a-Si based solar
cells. Intensive studies of solar cells with a graded i-layer showed a strong dependence
of the open circuit voltage (F^) and the fill factor (FF) on the cell designfl 10,111].
Generally the performance of a solar cell is characterized by the (F^) which is
the voltage supplied by the solar cell when the current is zero, the short circuit current
(Isc) which is defined as the current density if the voltage across the solar cell is zero,
and the fill factor (FF), which is defined as the ratio between the power density at the
maximum power point (Pmax), i.e. the point where the power supplied by the solar cell
is maximum, and the product of Voc and Isc. The optimum operating point on the I-V
characteristic curve of a solar cell can be found at which the output power (VmwcJmax) is
maximum. Thus the conversion efficiency (if) is defined by
Inputpower (Pin)
= "" "" xl00%
Where Vmax and Imax are the maximum values of the output voltage and output current.
The maximum power, the open circuit voltage Voc{Isc=Qi), and short circuit
current ISc(Voc=Q) are related as follows,
V x I x FF
'• 1 = ——7T
x 100%
0 2)
Here fill factor is given by
•""• - ^
J. Chevallier, H. Wieder, A. Onton and C. R. Guarnieri, Solid State Commun. 26
(1977) 867
[2] K. D. Mackenzie, J. R. Eggert, D. J. Leopold, Y. M. Li, S. Lin and W. Paul, Phys.
Rev. B 31 (1985)2198
K. D. Mackenzie, J. H. Burnett, J. R. Eggert, Y. M. Li and W. Paul, Phys.Rev. B
W. Paul, S. I. Iones, F. C. Marques, D. Pang, W. A. Turner, A. E. Wetsel and P.
Wickboldt, Mat. Res. Soc. Symp. Proc. 219 (1991) 211
M. Stutzmann, R. A. Street, C. C. Tsai, J. B. Boyce and S. E. Ready, J. Appl.
Phys. 66 (1989) 569
S. Aljishi, Z. E. Smith and S. Wagner, in Amorphous Silicon and Related
Materials, ed. by H. Fritzsche (World Scientific, Singapore, 1989) 887
[7] Y. Hishikawa, M. Sasaki, S. Tsuge, S. Okamoto and S. Tsuda, Proc. Conf. Of
Materials Research Society (San Francisco, 1993) 779
Y. Hishikawa, S. Tsuda, K. Wakisaka and Y. Kuwano, J. Appl. Phys. 73(1993)
[9] B. V. Roedern, D. K. Paul, J. Blake, R. W. Collins, G. Moddel and W. Paul, Phys.
Rev. B 25 (1982) 7678
[10] S. Nishikawa, H. Kakinuma, T. Watanabe and K. Nihei, Jpn. J. Appl. Phys. 25
A. H. Mahan, P. Raboisson and R. Tsu, Appl. Phys. Lett. 50 (1987) 335
[12] P. K. Banerjee, R. Dutta, S. S. Mitra and D. K. Paul, J. Non-Cryst. Solids 50
(1982) 1
[13] R. A. Rudder, J. W. Cook and G. Lucovsky, Appl. Phys. Lett. 43 (1983) 871
[14] S. Z. Weisz, Y. Goldstein, M. Gomez, J. A. Muir, O. Resto and R. Perez, Thin
Solid Films 119 (1984) 59
[15] N. Saito, K. Aoki, H. Sannomiya and T. Yamaguchi, Thin Solid Films 115
[16] N. Saito, T. Terada and T. Yamaguchi, J. Non-Cryst. Solids 83 (1986) 42
[17] W. Luft and Y. S. Tsuo, in Hydrogenated Amorphous Silicon Alloy Deposition
Processes (Dekker, New York, 1993)
[18] G. H. Bauer, C. E. Nebel and H. D. Mohring, Mater. Res. Soc. Symp. Proc. 118
(1988) 679
J. Yang, A. Banerjee, T. Glatfelder, K. Hoffinann, X. Xu and S. Guha,
Conference Record of the 24th IEEE Photovoltaic Specialists Conference
(Hawaii, 1994) 380
[20] J. Folsch, F. Finger, T. Kulessa, F. Siebke, W. Beyer and H. Wagner, Mater.
Res. Soc. Symp. Proc. 377 (1995) 517
[21] S. Wagner, V. Chu and J. P. Conde, 9th European Community Solar Energy
Conference (Freiburg, Germeny, 1989) 612
[22] G. Nakamura, U. Sato and Y. Yukimoto, Conference Record of the 16th IEEE
Photovoltaic Specialists Conference (San Diego, 1982) 1331
[23] A. Matsuda, K. Yagii, M. Koyama, M. Toyama, Y. Imanishi, N. Ikuchi and K.
Tanaka, Appl. Phys. Lett. 47 (1985) 1061
[24] A. Matsuda, M. Koyama, N. Ikuchi, Y. Imanishi and K. Tanaka, Jpn. J. Appl.
Phys. 25 (1986) L54
[25] T. Ichimura, T. Ihara, T. Hama, M. Ohsawa, H. Sakai and Y. Uchida, Jpn. J.
Appl. Phys. 25 (1986) L276
[26] K. Tanaka and A. Matsuda, Material Issues in Amorphous Semiconductor
Technology, ed. by A. Adler, Y. Hamakawa and A. Madan, Mater. Res. Soc.
Symp. Proc. 70 (1986) 245
[27] W. Paul, D. K. Paul, B. von Roedern, J. Blake and S. Ogur, Phys. Rev. Lett. 46
[28] I. Shimizu, M. Kumeda, A. Morimoto, Y. Tsujimura and I. Kobayashi, Material
Issues in Amorphous Semiconductor Technology, ed. by D. Adler, Y.
Hamakawa and A. Madan, Mater. Res. Soc. Symp. Proc. 70 (1986) 313
[29] K. Nozawa, Y. Yamaguchi, J. Hanna and I. Shimizu, J. Non-Cryst. Solids 59&60
[30] S. Oda, S. Takagi and I. Shimizu, Jpn. J. Appl. Phys. 25 (1986) 49
[31] N. Shibata, A. Tanabe, J. Hanna, S. Oda and I. Shimizu, Jpn. J. Appl. Phys. 25
(1986) L540
[32] S. Tsuda, H. Tarui, H. Haku, H. Dohjo, K. Watanabe, Y. Hishikawa, N.
Nakamura, Y. Nakashima, T. Takahama, H. Nishiwaki, T. Kukatsu, S.
Nakayama, M. Ohnishi, S. Nakano and Y. Kuwano, Jpn. J. Appl. Phys. 25
(1986) 1795
[33] S. Wagner, D. Slobodin, J. Kolodzey, S. Aljishi and R. Schwarz, in Amorphous
Metals and Semiconductors, ed. by P. Haasen and R. I. Jaffee (Pergamon,
Oxford, 1986) 334
[34] K. D. Mackenzie, J. Hanna, J. R. Eggert, Y. M. Li, Z. I. Sun and W. Paul, J. NonCryst. Solids 77&78 (1985) 881
[35] W. Paul, in Fundamental Physics of Amorphous Semiconductors, ed. by F.
Yonezawa (Springer, Heidelberg) 25 (1981) 72
[36] M. Zeman, M. J. Geerts, J. Siegl and J. W. Metselaar, Solar Energy Materials £3
(1991) 265
[37] W. Paul, in Amorphous Silicon and Related Materials, ed. by H. Fritzsche
(World Scientific, Singapore, 1989) 63
[38] S. Muramatsu, H. Kajiyama, H. Itoh, S. Matsubura and T. Shimada, in 20th IEEE
Photovoltaic Specialists Conf. Proc. (1988) 61
[39] G. H. Bauer, C. E. Nebel, M. B. Schubert and S. Schum, Mater. Res. Soc. Symp.
Proc. 149 (1989) 495
[40] A. Matsuda, J. Non-Cryst. Solids 59&60 (1983) 767
[41] S. Ray, A. R. Middya, S. C. De and A. K. Barua, Solar Energy Materials'23
[42] H. Haku, K. Sayama, E. Maruyama, H. Dohjoh, N. Nakamura, S. Tsuda, S.
Nakano, Y. Kishi and Y. Kuwano, Jpn. J. Appl. Phys. 30 (1991) 2700
[43] J. Xu, S. Miyazaki and M. Hirose, J. Non-Cryst. Solids 208 (1996) 277
[44] V. I.vKuznetsosgM. Zeman, L. L. A. Vosteen, B. S. Girwar and J. W. Metselaar,
J. Appl. Phys. 80 (1996) 6496
[45] D. A. Papaconstantopoulos and E. N. Economou, Phys. Rev. B 24 (1981) 7233
[46] L. Pauling, The Nature of the Chemical Bond (Cornell University Press, Ithaca,
New York, 1960) 3 rd ed. 78
[47] F. A. Cotton and G. Wilkinson, Advanced Inorganic Chemistry (Interscience,
New York, 1972) 3 rd ed. 113
[48] J. E. Huheey, Inorganic Chemistry (Harper and Row, New York, 1978) 2nd ed.
[49] A. Terakawa, M. Shima, K. Sayama, H. Tarui, S. Tsuda, H. Nishiwaki and S.
Nakano, Jpn. J. Appl. Phys. 32 (1993) 4894
[50] S. J. Jones, Y. Chen, D. L. Williamson, R. Zedlitz and G. Bauer, Appl. Phys.
Lett. 62 (1993) 3267
[51] B. Abeles and T. Tiedje, Phys. Rev. Lett. 51 (1983) 2003
[52] N. Ibaraki and H. Fritzsche, Phys. Rev. B 30 (1984) 5791
[53] L. Yang, B. Abeles and P. D. Persans, Appl. Phys. Lett. 49 (1986) 631
[54] J. Kakalios, H. Fritzsche, N. Ibaraki and S. Ovshinsky, J. Non-Cryst. Solids 66
[55] N. Maley and J. S. Lannin, Phys. Rev. B 31 (1985) 5577
[56] C. B. Roxlo, B. Abeles and P. D. Persans, J. Vac. Sci. Technol. B 4 (1986) 1430
[57] A. Asano, T. Ichimura, Y. Uchida and H. Sakai, J. Appl. Phys. 63 (1988) 2346
[58] K. Mui and F. W. Smith, J. Appl. Phys. 63 (1988) 475
[59] J. Kakalios and H. Fritzsche, Phys. Rev. Lett. 53 (1984) 1602
[60] S. C. Agarwal and S. Guha, Phys. Rev. B 31 (1985) 5547
[61] L. C. Tsai, C. F. Huang, J. C. Fan, Y. H. Chang and Y. F. Chen, J. Appl. Phys. 84
[62] P. V. Santos and W. B. Jackson, Phys. Rev. B 46 (1992) 4595
[63] H. C. Weller, S. M. Pasche, C. E. Nebel, F. Kessler, and G. H. Bauer, 19th IEEE
Photovoltaic Specialists Conf. Proc. (New Orleans, 1987) 872
[64] S. Guha, J. S. Payson, S.C. Agarwal and S. R. Ovshinsky, J. Non-Cryst. Solids
97&98 (1987) 1455
[65] R. A. Street, J. Non-Cryst. Solids, 137&138 (1991) 645
[66] M. Stutzmann, R. J. Nemanich and J. Stuke, Phys. Rev B 30 (1984) 3595
[67] D. Slobodin, S. Aljishi, R. Schwarz and S. Wagner, in Material Issues in
Applications of Amorphous Silicon Technology, ed. byD. Adler, A. Madan, and
M. J. Thompson, Materials Res. Soc. Symp. Proc. 49 (1985) 153
[68] H. Watanabe and K. C. Kao, J. Vac. Soc. Jpn. 24 (1981) 417
[69] P. A. Morin, N. W. Wang and S. Wagner, Mater. Res. Soc. Symp. Proc. 258
X. Xu, J. Yang and S. Guha, Appl. Phys. Lett. 62 (1993) 1399
[71] Q. Wang, H. Antoniadis, E. A. Schiff and S. Guha, Mater. Res. Soc. Symp. Proc.
[72] K. D. Mackenzie and W. Paul, Mater. Res. Soc. Symp. Proc. 95 (1987) 281
[73] T. Unold, J. D. Cohen and C. M. Fortmann, Mater. Res. Soc. Symp. Proc. 258
[74] M. Stutzmann, J. Nunnenkamp, M. S. Brandt and A. Asano, Phys. Rev. Lett. 67
[75] A. Rose, in Concepts in Photoconductivity and Applied Problems (Krieger, New
York, 1978)
[76] S. Wagner, V. Chu, J. P. Conde and J. Z. Liu, J. Non-Cryst. Solids 114 (1989)
[77] L. Tichy and A. Trisk, Solid State Commun. 41 (1982) 751
[78] L. Chen, J. Tauc, T. K. Lee and E. A. Stiff, J. Non-Cryst. Solids 114 (1989) 585
[79] E. Lotter and G. H. Bauer, J. Non-Cryst. Solids 114 (1989) 322
[80] G. Bruno, P. Capezzuto, G. Cicala and F. Cramarossa, J. Mater. Res. 4 (1989)
[81] T. Watanabe, K. Azumi, M. Nakatani and T. Shimada, Jpn. J. Appl. Phys. 29
(1990) 1419
[82] A. Morimoto, T. Miura, M. Kumeda and T. Shimizu, Jpn. J. Appl. Phys. 20
[83] W. Beyer, J. Non-Cryst. Solids 97&98 (1987) 1027
E. Maruyama, A. Terakawa, K. Sayama, K. Ninomiya, Y. Ffishikawa, H. Tarui,
S. Tsuda, S. Nakano and Y. Kuwano, 23rd IEEE Photovoltaic Specialists Conf.
(Louisville, 1993) 10-14
[85] D. E. Theodorou, H. J. Queisser and E. Bauser, Appl. Phys. Lett. 41 (1982) 628
[86] J. Y. Lin and H. X. Jiang, Phys. Rev. B 40 (1989) 10025
[87] M. I. Nathan, Solid-State Electron 29 (1986) 167
[88] H. J. Stromer, R. Dingle, A. C. Gossard, W. W. Wiegmann and M. D. Sturge,
Solid State Commun. 29 (1984) 705
[89] H. X. Jiang and J. Y. Lin, Phys. Rev. Lett. 64 (1990) 2547
[90] H. J. Queisser and D. E. Theodorou, Phys. Rev. Lett. 43 (1979) 401
[91] D. V. Lang and R. A. Logan, Phys. Rev. Lett. 39 (1977) 635
[92] M. K. Sheinkman and A. Y. Shik, Sov. Phys. Semicond. 10 (1976) 128
[93] H. J. Queisser and D. E. Theodorou, Phys. Rev. B 33 (1986) 4027
[94] D. E. Theodorou and H. J. Queisser, Appl. Phys. 23 (1980) 121
[95] H. J. Queisser, Phys. Rev. Lett. 54 (1985) 234.
[96] Y. Kuwano, M. Ohnishi, H. Shiwaki, S. Tsuda, T. Fukatsu, K. Enomoto, Y.
Nakashima and H. Tarui, Proc. 16th IEEE Photovoltaics Specialists Conf. (New
York, 1982) 1338.
[97] H. Sannomiya, S. Moriuchi, Y. Inoue, K. Nomoto, A. Yokota, M. Itoh, Y.
Nakata and T. Tsuji, Proc. 5th PVSEC (Kyoto, 1990) 387.
[98] S. Tsuda, T. Takahama, M. Isomura, H. Tarui, Y. Nakashima, Y. Hishikawa, N.
Nakamura, T. Matsuoka, H. Nishiwaki, S. Nakano, M. Ohnishi and Y. Kuwano,
Jpn. J. Appl. Phys. 26 (1987) 33.
[99] H. Tarui, T. Matsuyama, S. Okamoto, H. Dohjoh, Y. Hishikawa, N. Nakamura,
S. Tsuda, S. Nakano, M. Ohnishi and Y. Kuwano, Jpn. J. Appl. Phys. 28 (1989)
[100] S. Tsuge, Y. Hishikawa, N. Nakamura, S. Tsuda, S. Nakano, Y. Kishi and Y.
Kuwano, Tech. Dig. Int. Photovolt. Science and Engineering Conf.-5 (Kyoto,
[101] T. Sawada, H. Tarui, N. Terada, M. Tanaka, T. Takahama, S. Tsuda and S.
Nakano, Proc. 23rd IEEE Photovoltaic Specialists Conf. (Louisville, 1993).
[102] D. L. Staebler and C. R. Wronski, Appl. Phys. Lett. 31 (1977) 292.
[103] D. Adler, Solar Cells 9 (1983) 133.
[104] A. Terakawa, M. Shima, K. Sayama, H. Tarui, H. Nishiwaki and S. Tsuda, Jpn.
J. Appl. Phys. 34(1995)1741.
[105] S. Guha, J. Yang, A. Pawlikiewicz, T. Glatfelter, R. Ross and S. R Ovshinsky,
Appl. Phys. Lett. 54 (1989) 2330.
[106] Y. Nakata, H. Sannomiya, S. Moriuchi, A. Yokota, Y. Inoue, M. Itoh and H.
Itoh, Mater. Res. Soc. Symp. Proc. 192 (1990) 15.
[107] R. Araya, A. Catalano and S. R. Oswald, Appl. Phys. Lett 49 (1986) 1089.
[108] S. Guha, J. Yang, A. Pawlikiewicz, T. Glatfelter, R. Ross and S. R. Ovshinsky,
Conference Record of the 20th IEEE Photovoltaic Specialists Conference (Las
Vegas, 1988) 79.
[109] J. Yang, A. Banerjee and S. Guha, Mater. Res. Soc. Symp. Proc. 467 (1997)
[110] A. Pawlikiewicz and S. Guha, in ref. [108] p. 251.
[ I l l ] M. Block and D. Bonnet, 12th European Photovoltaic Solar Energy Conf.
(Amsterdam, the Netherlands, 1994) 1280.