UNITED STATES


 


 


Silicon-Based Third Generation Photovoltaics 

UMR CNRS 5270, INSA de Lyon, France


 


 


 


Article referenced as support for the following
sections:


 


 


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8


Silicon-Based Third Generation
Photovoltaics


Tetyana Nychyporuk and Mustapha Lemiti

University of Lyon, Nanotechnology
Institute of Lyon (INL),

UMR CNRS 5270, INSA de Lyon,

France


1. Introduction


In
order to ensure the widespread use of photovoltaic (PV) technology for
terrestrial applications, the cost per watt must be significantly lower than 1$
/ Watt level. Actually, the wafer based Silicon (Si) solar cells referred also
as the 1st generation solar cells are the most mature technology on
PV market. However such PV devices are material and energy intensive with
conversion efficiencies which do not exceed in average 16 %. In 2008 the
average cost of industrial 1 Wp Si solar cell with conversion efficiency of
14.5 % (multicrystalline Si cell of 150 x 150 mm2, 220 mm of thick,
SiN antireflecting coating with back surface field and screen printing
contacts) achieved approximately 2.1 € assuming the production volume of 30 -50
M Wp / per year (Sinke et al., 2008). At that cost level, the PV electricity
still remains more expensive comparing with traditional nuclear or thermal
power engineering. One of the most promising strategies for lowering PV costs
is the use of thin film technology, referred also as 2nd generation solar
cells. It involves low cost and low energy intensity deposition techniques of
PV material onto inexpensive large area low-cost substrates. Such processes can
bring costs down but because of the defects inherent in the lower quality
processing methods, have reduced efficiencies compared to the 1st generation
solar cells.


Material
limitations of the 1st generation solar cells and efficiency
limitations of the 2nd generation
solar cells are initiated boring of the Si-based 3rd generation
photovoltaic. Its main goal is to significantly increase the conversion
efficiency of low-cost photovoltaic product. Indeed, the Carnot limit on the
conversion of sunlight to electricity is 95% as opposed to the theoretical
upper limit of 30% for a standard solar cell (Shockley & Queisser 1961).
This suggests the performance of solar cells could be improved 2 - 3 times if
different concepts permitting to reduce the power losses were used.


The
two most important power loss mechanisms in single-band gap photovoltaic cells
are (1) the inability to absorb
photons with energy less than the band gap and (2)
thermalisation of photon energy exceeding the gap (Fig. 1). Longer
wavelength is not absorbed by the solar cell material. Shorter wavelength
generates an electron-hole pair greater than the bandgap of the p-n junction
material. The excess of energy is lost as heat because the electron (hole)
relaxes to the conduction (valence) band edges. The amounts of the losses are
around 23 % and 33 % of the incoming solar energy, respectively (Nelson, 2003).
Other losses are junction loss, contact loss and the recombination loss. Theory
predicts (Shockley & Queisser 1961) that the highest single - junction
solar cell efficiency is roughly 30%, assuming such factors as the intensity of
one sun (no sunlight concentration), a one junction solar cell (a single
material with a single bandgap), and one electron-hole pair produced from each
incoming photon.











	
  

 	
  

 
	
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To
efficiently convert the whole solar spectrum into the electricity three main families
of approaches have been proposed (Green et al., 2005) (Green, 2002): (i) increasing the number of bandgaps (tandem cell concept); (ii) capturing carriers before
thermalisation, and (iii) multiple
carrier pair generation per high energy photon or single carrier pair
generation with multiple low energy photons. Up to now, tandem or in other
words multijunction cells provide the best-known example of such
high-efficiency approaches. Indeed, the loss process (2) of Fig. 1 can be largely eliminated if the energy of the
absorbed photon is just a little higher that the cell bandgap. The concept of
tandem solar cells is based on the use of several solar cells (or subcells) of
different bandgaps stacked on top of each other (Fig. 2), with the highest
bandgap cell uppermost and lowest on the bottom. The incident light is
automatically filtered as it passes through the stack. Each cell absorbs the
light that it can most efficiently convert, with the rest passing through to
underlying lower bandgap cells (Green et al., 2007). The using of multiple
subcells in the tandem cell structure permits to divide the broad solar
spectrum on smaller sections, each of which can be converted to electricity
more efficiently. Performance increases as the number of subcells increases,
with the direct sunlight conversion efficiency of 86.8 % calculated for an
infinite stack of independently operated subcells (Marti & Araujo, 1996).
The efficiency limit reaches 42.5 % and 47.5 % for 2- and 3-subcell tandem
solar cells (Nelson, 2003) as compared to 30% of one junction solar cell.





Fig.
1. Loss processes in a standard solar cell: (1)
non-absorption of below band gap photons; (2) lattice thermalisation loss; (3) and (4) junction
and contact voltage losses; (5) recombination
loss (Green, 2003).


 





Fig. 2. Tandem cell approach (Green, 2003).


Having
to independently operate each subcell is a complication best avoided. Usually,
subcells are designed with their current output matched so that they can be
connected in series. This constrain reduces performance. Moreover, it makes the
design very sensitive to the spectral content of the sunlight. Once the output
current of one subcell in a series connection drops more than about 5% below
that of the next worst, the best for overall performance is to short-circuit
the low-output subcell, otherwise it will consume, rather than generate power.


It
should be also noted the common point of confusion about solar cells
efficiency. The measured efficiency of solar cell depends on the spectrum of
its light source. The space solar









	
  

 	
  

 
	
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spectrum
or air mass zero (AM0) spectrum is richer in ultraviolet light than the typical
terrestrial solar spectrum (air mass 1.5 or AM1.5). Taking into account that
the ultraviolet light is converted into electricity less efficiently than the
other parts of the spectrum, the resulting efficiencies for AM0 are thus lower
(Green et al., 2010). Since cells are typically measured under the spectrum for
their intended use and efficiencies are not easily converted, this chapter will
indicate efficiencies measured under non-concentrated AM1.5 at 250 unless
otherwise specified.





Fig. 3. Schematic view of GaInP/GaAs/Ge solar cell.





Fig. 4. Using of multijunction solar cells for Mars rover missions 1.


Up
today the tandem cells have been developing on monolithic integration of
non-abundant III –V materials by means of rather expensive technologies of
fabrication like molecular beam epitaxy (MBE) or metal-organic chemical vapor
deposition (MOCVD). Currently commercially available multijunction cells
consist of three subcells (GaInP/GaAs/Ge), which all have the same lattice
parameter and are grown in a monolithic stack (Fig. 3). The subcells in this
monolithic stack are series connected through the tunnel junctions. The record
efficiency of 32% was achieved in 2010 for this type of cells (Green et al.,
2010). These high-efficiency solar cells are being increasingly used in solar
concentrator systems, where development of both the solar cells and the
associated optical and thermal control elements are actively being pursued. The
performance of tandem solar cells has been demonstrated, but work is continuing
to increase the numbers of junctions and optimize the bandgap junctions. The
choice of materials with optimal or near-optimal bandgap is severely limited by
the lattice matching constraint of these cells. Another approach to increasing
of multijunction solar cell efficiency is the incorporation of materials with a
mismatch in the lattice constant. Graded composition buffers between the
lattice mismatched subcells are used to reduce the density of the threaded
dislocations resulting from the lattice mismatch strain. Lattice mismatch
technology opens the parameter space for junction materials, allowing the
choice of materials with more optimal bandgaps and a potential for higher cell
efficiency.


	
  

 
	
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 1 http://marsrovers.nasa.gov/gallery/press/spirit/20060104a.html

 













	
  

 	
  

 
	
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In
2010 the cost of multijunction solar cells still remains too high to allow
their use outside of specified applications (for example space applications,
Mars rover missions (Crisp et al., 2004) (Fig. 4)...). The high cost is mainly
due to the complex structure and the high price of materials. In this context
the fabrication of multijunction solar cells on the base of abundant low-cost
materials that do not cause toxicity in the environment and by using approaches
amenable to large scale mass production, like thin film deposition techniques,
remains challenging.


2. Silicon based tandem cells


Silicon
is a benign readily available material, which is widely used for solar cell
fabrication. It has a bandgap of 1.12 eV at 300 K, which is close to optimal
not only for standard, single p – n junction cell, but also for the bottom cell
in a 2-cell or even a 3-cell tandem stack (Conibeer et al., 2008). Therefore a
solar cell entirely based of Si and its dielectric compounds (referred also as all-Si tandem solar cell) with other
abundant elements (i. e. silicon dioxide, nitrides or carbides) fabricated with
thin film techniques, is advantageous in terms of potential for large scale
manufacturing and in long term availability of its constituents. As was already
mentioned previously, thin film low-temperature deposition techniques results
in high defect density films. Hence solar cells must be thin enough to limit
recombination due to their short diffusion lengths, which in turn means they
must have high absorption coefficients.


For
AM1.5 solar spectrum the optimal bandgap of the top cell required to maximize
conversion efficiency is 1.7 to 1.8 eV for a 2-cell tandem with a Si bottom
cell and 1.5 eV and 2.0 eV for the middle and upper cells for a 3-cell tandem
(Meillaud et al., 2006). It should be also noted that for terrestrial
applications (AM1.5 solar spectrum), the highest bandgap necessary for the
Si-based tandem solar cells is limited to 3.1 eV, the energy at which the
absorption from the encapsulation material, such as ethylene-vinyl acetate
(EVA), starts to play an important role.


2.1 Quantum confinement in Si
nanostructures


To
increase Si bangap, nanoscale size dependent quantum confinement effect can be
used. Indeed, the quantum confinement effect manifests itself by significant
modification of electronic band structure of Si nanocrystals when their size is
reduced to below the exciton Bohr radius (4.9 nm) of bulk Si crystals. In
particular, quantum confinement effect provokes the increasing of the effective
bandgap of Si nanocrystals. Moreover, for indirect bandgap semiconductors, like
Si, geometrical confinement of carriers increases the overlap of electron and
hole wavefunctions in momentum space and thus enhances the oscillator strength
and as a consequence increases its absorption coefficient. From this effect,
one can expect Si nanocrystals to behave as direct bandgap semiconductors.
However, there is some evidence suggesting that the momentum conservation rule
is only partially broken and Si nanocrystal strongly preserves the indirect
bandgap nature of bulk Si crystals (Kovalev et al., 1999). Si nanostructures
are thus the perfect candidates for higher bandgap materials in all-Si tandem
cell approach.


Since
the observation in 1990 of strong room-temperature photoluminescence from
nanostructured porous Si (Canham, 1990), significant scientific interest has
been focused of the simulation of optical and electrical properties of Si
nanostructures regarding their size, shape, surface termination, number and
degree of interconnections, impurity doping and so









	
  

 	
  

 
	
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on.
Many reviews addressing this problematic since appeared (one of the good recent
reviews is Ref (Bulutay & Ossicini, 2010)). In this paragraph we will only
underline some important points resulting from quantum confinement effect.


Ab initio calculations using density functional theory (DFT) indicate that the
increasing of the optical bandgap of Si nanocrystals (or in other words quantum
dots (QDs)) is expected to vary from 1.4 to 2.4 eV for a nanocrystal size of
8-2.5 nm (Ogut et al., 1997). However, further DFT calculations have found that
in addition to quantum confinement effect in small QDs, the matrix has a strong
influence of the resulting energy levels (KÖnig et al., 2009). With increasing
polarity of the bonds between the nanocrystal and the matrix, there is an
increasing dominance of the interface strain over quantum confinement. For a 2
nm diameter nanocrystal, this strain is such that the highest occupied
molecular orbital (HOMO)-lowest unoccupied molecular orbital (LUMO) gap is
significantly reduced in a polar SiO2 matrix but not much affected
in a less polar SiNx or nonpolar SiC matrix (KÖnig et al., 2009). It
is also shown the reduction in gap energy on going from a QD in vacuum to the
one embedded in a dielectric. An additional freedom of material design can be
also introduced by impurity doping and interconnections between Si QDs, the
both ones modify its optical and electrical transport properties (Nychyporuk et
al., 2009) (Mimura et al., 1999).


3. Silicon quantum dot solar cells


Si
QDs offer the potential to tune the effective bandgap, through quantum
confinement, and allow fabrication of optimized tandem devices in one growth
run in a thin film process.


3.1 Fabrication of Si QD
nanostructures


Different
technological approaches allowing formation of Si quantum dots in a dielectric
matrix have already been developed, permitting to obtain Si nanocrystals as
small as 1 nm in diameter. Between the most used deposition techniques one can
cite reactive evaporation, ion implantation, sputtering and plasma enhanced
chemical vapor deposition (PECVD). Considerable work has been done on the
growth of Si nanocrystals embedded in silicon oxide dielectric matrices (SiO2)
(Zacharias et al., 2002) (Stegemann et al., 2010), silicon nitride (Si3N4)
(Kim et al., 2006) (Cho et al., 2005) (Mercaldo et al., 2010) (So et al., 2010)
and silicon carbide (SiC) (Song et al., 2008) (Kurokawa et al., 2006) (Gradmann
et al., 2010) (LÖper et al., 2010) (Cho et al., 2007). An accurate control of
the size and density of Si nanocrystals is mandatory in bandgap engineering for
solar cell applications. It should be noted that Si nanocrystals prepared by
different methods present slightly different properties which depend on the
preparation procedure, because of different defect density, different degree of
interconnections between the nanocrystals, different surface termination, and
so on. Conventionally, Si QDs in dielectric matrix like SiO2, Si3N4
and SiC can be synthetized by self-organized growth from Si rich dielectric
layers, which are thermodynamically unstable and therefore undergo phase
separation upon appropriate post-annealing step to form nanocrystals. For
example for Si rich oxide layer, the precipitation occurs according to the
following:





It
should be noted that due to the fact that both the polarity and length of the
bonds decrease towards those of Si-Si for SiO2 to Si3N4
and SiC, this implies that the segregation









	
  

 	
  

 
	
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and
precipitation effect for Si in the three matrices would decrease such that in
SiC formation of QDs is likely to be most difficult.





(a)





(b)


Fig.
5. (a) Multi-layer structure illustrating precipitation of Si QDs in a Si-rich
layer; (b) TEM image of a superlattice of Si QDs in SiO2 matrix
(Conibeer et al., 2008).


The
QD size and density control are normally realized by changing the chemical
stoichiometry of the bulk films. By reducing the Si-richness in a bulk Si-based
matrix, smaller nanocrystals can be achieved. Nevertheless, this will
simultaneously reduce the density of nanocrystals in the film due to the
reduced Si-richness. The low density of QDs with desired size leads to a
negative effect on the electrical conductivity of the films. Moreover, the
assumption that all the excess Si precipitates to nanocrystals turns out to be
oversimplification. In fact, it has been observed that only half the excess Si
clusters in these precipitates upon annealing at 10000C for 30 min, in material
deposited by PECVD, with a considerable amount of suboxide material forming in
the matrix. Therefore, the method allowing a fabrication of high density, but
narrow size distributed Si QDs films via superlattice approach, firstly
reported by Zacharias (Zacharias et al., 2002) , was adopted by the majority of
researchers. It should be mentioned that even without using the multilayer






	
  

 	
  

 
	
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approach
it is still possible to fabricate well – ordered and uniform Si nanocrystals in
a film with a high dot density (Surana et al., 2010).


3.1.1 Si QDs in silicon oxide matrix


A
simple technique to prepare the multilayer structure known also as
superlattices of Si QDs in silicon oxide matrix was firstly reported by
Zacharias (Zacharias et al., 2002) . It consists in deposition of alternating
layers of stoichometric Si oxide (SiO2) and Si rich oxide (SRO) of
thicknesses down to 2 nm. This precision is normally achieved by using RF
magnetron sputtering or plasma enhanced chemical vapor deposition (PECVD)
technique. The deposition consisting typically of 20-50 bi-layers is followed
by the annealing step is N2 ambient from 1050 to 11500C for 1 h.
During the annealing step, the surface energy minimization favors the
precipitation of Si in the SRO layer into approximately spherical QDs (Conibeer
et al., 2008). This process is illustrated on Fig. 5 (a). The diameter of Si
QDs is constrained by the SRO layer thickness and quite uniform size dispersion
is achieved within about 10% (Zacharias et al., 2002) . The density of the QDs
can be varied by the composition of the SRO layer. Fig. 5 (b) shows typical
transmission electron microscope (TEM) of the multi-structure SiQDs in SiO2
matrix grown by this method. TEM evidence indicates that these nanocrystals
tend to be spherical – as surface energy minimization would dictate - and at
this scale would have energy levels confined in all three dimensions and hence
can be considered as quantum dots.


Nowadays,
the phase separation, solid state crystallization and optical properties of SiO2/SRO/SiO2
superlattices are already well understood. However, it is a major challenge to
achieve charge carrier transport through a network of Si QDs embedded in a SiO2
matrix. Therefore, other Si based host matrices such as Si3N4
or SiC that feature lower energy band offsets with respect to the Si band edges
and thus higher carrier mobility are attractive.


3.1.2 Si QDs in silicon nitride matrix


For
the reasons stated above, it was explored the fabrication of Si QDs in silicon
nitride matrix. Thick layers of silicon-rich nitride, when annealed at above
10000C, precipitate to Si QD (Kim et al., 2005). Multilayered structures also
result in Si QD formation with controlled size of the Si QDs (Cho et al.,
2005). The annealing temperature can also be used to modify the nitride matrix,
with it being amorphous below 11500C but with crystalline nitride phases, in
addition to the Si QDs, appearing at temperatures ranging from 1150 to 12000
(Scardera et al., 2008). Multilayered structures can be deposited by sputtering
or by PECVD with growth parameters and annealing conditions very similar to
those for oxide giving good control of QD sizes. The main difference is the
extra H incorporation with PECVD that requires an initial low-temperature
anneal to drive off excess hydrogen and prevent bubble formation during the
high-temperature anneal (Cho et al., 2005). Si QDs can also be grown in situ during PECVD deposition, where
they form in the gas phase (Lelièvre et al., 2006). There is much less control
over size and shape but no high-temperature anneal is required to form the Si
QDs (Fig. 6 (a)). Multilayer growth using this in-situ
technique has also been attempted with irregular shaped but
reasonably uniform sized Si QDs (Fig. 6 (b)).


The
formation of Si quantum dots in SiO2/Si3N4
hybrid matrix was also reported (Di et al., 2010). In this approach alternating
silicon rich oxide and Si3N4 layers were produced
followed by post-deposition anneals. In addition, it should be noted that Si3N4
acts as a better diffusion barrier compared to SiO2. It restricts
the displacement of Si atoms as well as dopant atoms under high processing
temperatures.









	
  

 	
  

 
	
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Fig.
6. In-situ grown Si QDs in the
gas phase and dispersed in Si3N4 matrix: (a) One layer
structure (Lelièvre et al., 2006); (b) multi-layer structure (Conibeer et al.,
2008).


3.1.3 Si QDs in silicon carbide
matrix


Si
QDs in a SiC matrix offer an even lower barrier height and hence potentially
better electronic transport properties. However, the low barrier height also
limits the minimum size of QDs to about 3 nm or else the quantum-confined
levels are likely to rise above the level of the barrier, which should be
around 2.3 eV for amorphous SiC. Si QDs in SiC matrix have been formed in a
single thick layer by Si-rich carbide deposition followed by high-temperature
annealing at between 8000 and 11000C in a very similar process to that for oxide
(Fig. 7). Si1-xCx/SiC multilayers have also been
deposited by sputtering to give better control over the Si QD as with oxide and
nitride matrices. However, contrary to these previous matrices, the both Si and
SiC QDs have been produced by high temperature annealing of Si –rich SiC layer
or in a SiC1-xCx/SiC multistructure. The formation of SiC
nanocrystals can hinder the formation of Si QDs.





Fig. 7. Cross-sectional HRTEM image of Si-rich SiC layer after the
thermal annealing (Conibeer, 2010).


 





Fig. 8. TEM image interconnected Si QDs forming thin films.









	
  

 	
  

 
	
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3.1.4 Interconnected Si QDs
forming thin films


To
be successfully applied as a material for all-Si tandem solar cells, the small
size of Si QDs is not the single prerequisite. It is also necessary to assure
their high density in order to achieve a direct tunneling of the photogenerated
charge carriers between the QDs. This still constitutes the bottleneck of the
approaches cited above. Recently, the fabrication of thin films composed by
highly packed Si QDs with a controlled bandgap values was reported (Nychyporuk
et al., 2009). This approach is based on the in-situ
nucleation of Si QDs in the gas phase during PECVD deposition by
using SiH4 as a gas precursor. Indeed, the dust particle formation
in Ar-SiH4 plasma is known to be a time-dependent four step process
occurring in the gas phase: (i) polymerization
phase, (ii) accumulation phase, (ii) coalescence and (iv) surface deposition growth. During the
polymerization phase, the nucleation of extremely small particles (1 nm) takes
place. They progressively grow in size with time and at the end of the
polymerization phase, starting from about 1 nm, a short accumulation phase
begins. During this phase the nanoparticles size remains constant and only
their density increases in the plasma environment. The coalescence phase starts
once the nanoparticles critical density is reached. The small nanoparticles (1
nm) begin to agglomerate at least two by two to form larger nanoclusters. In
consequence, a number of interconnections between the nanoparticles increases.
The final phase corresponds to the plasma species deposition on the surface of
strongly agglomerated nanoparticles. During this phase a hydrogenated amorphous
Si shell layer is formed around the crystalline Si core. The thickness of this
amorphous shell increases with time. The square wave modulation of the power
amplitude applied to the plasma has been found as a suitable technique
permitting to obtain the deposition of Si QDs with required size. It consists
of alternating periods of plasma switching time followed by the plasma
extinction time. As a result, Si QDs grown in the gas phase during the plasma
switching time were deposited on a substrate (Fig. 8). The careful tuning of
the plasma switching time permits to precisely control the phase of Si QD
growth and as a consequence their size and degree of interconnections between
them. Si QD based thin films deposited under dusty plasma conditions appear to
be promising candidates for all-Si tandem solar cell applications.


3.2 Shallow-impurity doped Si
nanostructures


A
requirement for a tandem cell element is the presence of some form of junction
for carrier separation. Phosphorous (P) and boron (B) are excellent dopants in
bulk Si as they have a high solid solubility and alter the conductivity of the
bulk Si by several orders of magnitude. Hence they are good initial choices to
study the doping in the Si nanocrystals. Doping of Si nanostructures is a subject
of intense research (Tsu et al., 1994) (Holtz & Zhao, 2004) (Erwin et al.,
2005) (Norris et al., 2008) (Ossicini et al., 2006). Unfortunately, the main
difficult in existing doping techniques arises from the fluctuation of impurity
number per nanocrystal in a nanocrystal assembly. For Si nanocrystals as small
as few nanometers in diameter, the expression of the doping level in the form
of “impurity concentration” is not suitable and it should be expressed as
“impurity numbers” because it changes digitally. For example, doping of one
impurity atom into a nanocrystal of 3 nm in diameter ( about 700 atoms)
corresponds to an impurity concentration of 7.0 X 1019 atoms/cm3.
At this doping level, bulk Si is a degenerate semiconductor and exhibits
metallic behavior. However, by means of electron spin resonance (ESR)
spectroscopy it was shown that Si nanocrystals do not become metallic even
under heavily doped conditions. Therefore, in nanocrystals, addition or
subtraction of a single impurity atom drastically changes the electronic
structure









	
  

 	
  

 
	
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and
the resultant optical and electrical transport properties. So, for solar cell
application, the development of a technique permitting to control the “impurity
number” with extremely high accuracy is indispensable.


Up
to now, an accurate control of “impurity number” in a Si nanocrystal has not
been achieved. One of the largest problems of the growth of doped Si
nanocrystals is that impurity atoms are pushed out of nanocrystals to
surrounding matrices by the so-called self-purification effect. This effect can
be understood by considering very high formation energy of doped nanocrystals
(Ossicini et al., 2005). Impurity concentration in nanocrystals is thus always
different from average concentration in a whole system. In the worst case, the
number of impurity in a nanocrystal becomes zero even when average
concentration is rather high. The development of viable technique to
characterize impurities, especially “active” impurities doped into
nanocrystals, is crucial. The resistivity measurements are thus complemented
with ESR spectroscopy as well as PL spectroscopy.


As
it was discussed previously, numerous methods have been reported for the growth
of intrinsic Si QDs. On the other hand, a limited number of studies are
published concerning the growth of shallow impurity-doped Si QDs with the
diameter below 10 nm. One of the mostly used methods for shallow-impurity
doping of Si nanocrystals is plasma decomposition of SiH4 by adding
dopant precursors (diborane (B2H6) and phosphine (PH3))
(Pi et al., 2008) (Stegner et al., 2008). This method permits to obtain a
variety of morphologies from densely packed nanocrystalline films to
nanoparticle powder by controlling process parameters (Nychyporuk et al.,
2009). Another method is incorporation of doping atoms into SRO layers by
simultaneous co-sputtering of Si, SiO2 and P2O5
(or B2O3) in SiO2/SRO superlattice approach
described previously (Mimura et al., 2000). During the annealing, Si
nanocrystals are grown in phosphosilicate (PSG) (n-type Si QDs) or borosilicate
(BSG) (p-type Si QDs) thin films. It should be also noted that the impurity
concentration in nanocrystals is different from that of the matrices because
the segregation coefficient strongly depends on the kind of impurities and
surrounding medium.





Fig.
9. PL spectra of (a) B-doped (Mimura et al., 1999) and (b) P-doped Si
nanocrystals (Fujii et al., 2002) at room temperature for different doping
concentrations.


The
presence of impurity atoms inside Si nanocrystals can be confirmed by the PL
spectroscopy. Indeed, the introduction of extra carriers by impurity doping
makes the three-body Auger process possible (Kovalev et al., 2008). In Auger
recombination, the energy of









	
  

 	
  

 
	
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an
electron-hole pair is not released in the form of photon. This energy is given
to a third carrier, which after the interaction losses its excess energy as
thermal vibrations. Since this process is a three-particle interaction, it is
normally only significant in strongly non-equilibrium conditions or when the
carrier density is very high. For doped nanocrystals the value of Auger rate is
four to five orders of magnitude larger than the radiative rate of excitons,
and thus one shallow impurity can almost completely kill PL from the
nanocrystal. Hence with increasing of average impurity numbers in a nanocrystal
assembly, the PL intensity is expected to decrease. This effect was really
observed in p-type Si nanocrystals. As one can see on Fig. 9 (a) with
increasing of B concentration, the PL intensity monotonously decreases (Fujii
et al., 1998) (MÜller et al., 1999) (Stegner et al., 2008).


In
P-doped Si nanocrystals, the situation is different. When the phosphorous
concentration is relatively low, the PL intensity increases slightly compared
to that of the undoped Si nanocrystals (Fig. 9 b) (Fujii et al., 2000) (Mimura
et al., 2000) (Tchebotareva et al., 2005). The increase of the PL intensity
indicates that non-radiative recombination processes are quenched by P doping.
One of the possible explanation is that electrons supplied by P are captured by
the dangling bonds, which inactivate the nonradiative recombination centers and
compensate donors (Stegner et al., 2008) (Lenahan et al., 1998). It should be
also noted that the PL intensity also strongly depends on the size of the
shallow-doped Si nanocrystals. There are many theoretical studies on
preferential localization of impurities in Si nanocrystals. It should to be
noted that it is almost impossible to control experimentally the location of
impurities in nanocrystals. However, the information on localization was
experimentally obtained (Kovalev et al., 1998). It was shown that P dopants are
localized at or close to the surface of Si nanocrystal. On the contrary, the B
atoms are primary incorporated into the Si nanocrystal core. However, the
preferential localization of impurities may depend on nanocrystal growth
process and the surface termination. Therefore, properties of dopant may be
quite different between Si nanocrystals grown by the decomposition of SiH4,
phase separation of SRO and so on. In any cases the doping efficiency by B
atoms is much smaller than that of P atoms due to larger formation energy of
B-doped Si nanocrystals than P-doped ones (Kovalev et al., 1998) (Ossicini et
al., 2005). To what concerns the optical and electrical properties, contrary to
the intrinsic Si nanocrystals, the shallow doped Si nanocrystals present new
degree of freedom to control them. For example, the size is one of the main
parameters to control optical bandgap of the intrinsic Si nanocrystals. On the
other hand, due to the difference in the electronic band structure in the case
of doped and codoped Si nanocrystals (obtained by the simultaneous doping by B
and P atoms), the optical bandgap is determined by the combination of the size
and impurity concentration (Fujii et al., 2010).


It
is worth noting that the impurity atoms alter the formation kinetics of Si
nanocrystals. Indeed, the average size of P-doped Si nanocrystals is increased
compared to the undoped ones under the same experimental conditions (Conibeer
et al., 2010), and in some experiments this increasing was almost double.
Contrary to the doping with P, the doping with B results in the forming of
smaller Si nanocrystals compared to the undoped case (Hao et al., 2009). The
crystalline volume fraction was found to decrease with increasing of B
concentration (Hao et al., 2009), which suggests that boron suppresses Si
crystallization. One of the possible reasons is the local deformations induced
by the impurity atoms.


3.2.1 Dark resistivity
measurements


The
resistivity of the Si QD material is an important parameter for photovoltaic
applications. The influence of the doping concentration on resistivity of Si QD
superlattices was studied









	
  

 	
  

 
	
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(Hao
et al., 2009, 2009a), (Conibeer et al, 2010) (Ficcadenti et al, 2009). The
contact resistances in the above measurements were determined by using the TLM
(Transmission Line Model) method proposed by Reeves and Harrison (Reeves &
Harrison, 1982). This method involves measurement of the resistance between
several pairs of contacts, which have identical areas, but are separated by
different spaces (Fig. 10). The dark resistivity is then defined as:Pdark =V X d X w/I X l , where d is the thickness of Si QD superlattice, w is the length of Al pad and l is the spacing of Al pads.


To
perform the dark resistivity measurements, the Si QD superlattices were grown
on the quartz substrate. The ohmic contacts were obtained by thermal
evaporation of Al, followed by sintering at a temperature (500-5300C) lower
than the Al-Si eutectic temperature to allow the Al to spike down into the film
(Voz et al, 2000). The schematic view of the final structure for the lateral
resistivity measurements is presented on the Fig. 10.


Fig.
11 (a) and (c) represents the room temperature dark resistivity of Si QD/SiO2
multilayer films for various phosphorous and boron doping levels, respectively.
As one can see, the introduction of a slight amount of P and B drastically
changes the dark resistivity of the films, from 108 Wcm for the
undoped samples to 102 - 10 Wcm for doped ones, which is 6 -7 orders
of magnitude lower than that of the undoped samples. This decrease in
resistivity may be the consequence of an increase in mobile carrier
concentration due to a rise in the number of active dopants in the film.


The
TLM method was also used to measure the temperature dependence of the
resistance of the Si QD films with various (b) phosphorous (Hao et al, 2009)
and (d) boron (Hao et al., 2009) concentrations (Fig. 11 (b) and (d),
respectively). These measurements permit to estimate the values of the
activation energy (Ea),
that is in a n- (p-) doped semiconductor the energy difference between the
conduction (valance) band and Fermi level. The activation energy was calculated
by using relation R " exp(Ea/kT). As one can see,
with the increasing of the doping level, for both types of impurities the
activation energy decreases from 0.5 eV to 0.1 eV. This result is consistent
with the view that the observed resistivity decreases are a consequence of an
increase in carrier concentration due to more active dopants in the film. The
decrease in Ea accompanying
the drop in resistivity indicates that the Fermi level energy is moving toward
the conduction (valance) band for n- (p-) type doped Si QDs.





Fig. 10. Schematic layout of the Al contacts on a film for dark
resistivity measurements (Hao et al., 2009).









	
  

 	
  

 
	
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Fig.
11. Dark resistivity of Si QD/SiO2 multilayer films for various (a)
phosphorous (Hao et al., 2009) and (c) boron (Conibeer et al., 2010) doping
levels.; Temperature dependence of the resistance of the Si QD films with
various (b) phosphorous (Hao et al., 2009) and (d) boron (Hao et al., 2009)
concentrations .


3.3 Optical properties of Si QDs


Regarding
to photovoltaic applications, the optical bandgap and the absorption
coefficient of Si QDs are the very first physical parameters to be studied and
optimized prior to solar cell fabrication. Optical methods provide an easy and sensitive
tool for measuring the electronic structure of quantum objects, since they
require minimal sample preparation and the measurements are sensitive to the
quantum effects. The energy gaps of Si QDs could be determined, for example,
from photoluminescence (PL) measurements, whereas the absorption coefficient
from the transmission-reflection measurements.


3.3.1 Bandgap of Si QDs


Experimental
energy gaps of isolated Si QDs in SiO2 and SiNx matrices
reported by several research groups are shown on Fig. 12 (Cho et al., 2004)
(Park et al., 2000) (Takeoka









	
  

 	
  

 
	
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et
al., 2000). As one can see the bandgap values of Si nanostructured material
could be adjusted in the large range (up to 3.1 eV), covering an important part
of the solar spectrum. The results obtained by different teams are in good
agreement where the matrix is the same but are quite different for QDs in oxide
compared to nitride, particularly for small QDs. They are also qualitatively
consistent with the results from ab-initio modeling
(KÖnig et al., 2009) (Ogut et al., 1997), which had been carried out for the
confined energy levels in Si QD consisting of a few hundred atoms. One can
observe the expected increasing of confinement energy with decreasing QD size,
but also that the amino-terminated QDs (silicon nitride) have energies about
0.5 eV more than the hydroxyl-terminated ones (silicon oxide). The last one
observation is consistent with the explanation for the enhanced energies of QDs
in nitride given by Yang et al (Yang et al., 2004), that the reason for it is
due to better passivation of Si QDs by nitrogen atoms eliminating the strain at
the Si/Si3N4 interface.


Influence
of interconnections between the QDs on tuning of their bandgap was also studied
(Nychyporuk et al., 2009) (Degoli et al., 2000). Indeed, electronic coupling
between the neighboring low-dimensional Si nano-objects constituting a complex
quantum system must be considered. This coupling leading to intense energy
transfer processes between the electronically communicating quantum objects
determines physical properties of the whole quantum system and, therefore, has
to be taken into account, of course. It was shown that when the nanocrystals
(Allan & Delerue, 2007) (Bulutay, 2007) start to touch each other, the
bandgap value of the assembly begins to decrease rapidly (Fig. 13). Thus, the
bandgap of the interconnected nanostructures depends not only on the
nanocrystal dimension but also on the degree and number of interconnections between
them.





Fig.
12. Experimental energy gaps of three-dimensionally confined Si QDs in SiO2
and SiNx matrices (3000C) for several research groups (Takeoka et
al., 2000) (Kim et al., 2004), (Kim et al., 2005) (Yang et al., 2004)
(Fangsuwannarak, 2007).





Fig.
13. Evolution of the bandgap of the quantum system constituted of 27
interconnected Si QDs as a function of the distance between the QDs.






	
  

 	
  

 
	
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3.3.2 Optical absorption of Si QDs


The
absorption coefficient was experimentally determined for Si QDs in deferent
matrices. Fig. 14 shows the global absorption coefficient of SiNx
layers of different stoichiometries with Si QDs embedded inside (Nychyporuk et
al., 2008). The absorption coefficient of polycrystalline silicon (poly-Si) is
also added for comparison. As it can be seen, the global absorption coefficient
of the composite SiNx decreases with stoichiometric ratio R (i.e.
with decreasing of Si QD size and density) and its band-edge shifts to higher
energies. Its magnitude is being much lower than that one of the absorption
coefficient of poly-Si. No evidence of oscillator strength enhancement was
observed and the global absorption coefficient is limited principally by the
volume fraction of Si QDs inside the dielectric matrices. These why, the
maximum absorption coefficient, approaching this one of the bulk Si, was found
in the case of interconnected Si QDs forming thin films (Nychyporuk et al.,
2009).





Fig.
14. Global absorption coefficient of SiNx layers of different
stoichiometries with embedded Si QDs. The absorption coefficient of poly-Si is
also presented for comparison (Nychyporuk et al., 2008).


3.4 Electrical transport
mechanisms in Si QD ensembles


While
the optical properties of the various ensembles of individual Si nanocrystals
have been investigated by many researchers, relatively little attention was
paid to the transport properties of 3D ensembles of such QDs. In this paragraph
we will only briefly review the main results on the transport mechanisms
obtained previously in the literature. A complete review of the electrical
transport mechanisms in 3D ensembles of disordered Si nanocrystallites embedded
in insulating continuous matrices can be found in Ref. (Balberg et al., 2010).
To what concerns the transport processes in nanocrystalline Si superlattices,
they were well reviewed in Ref. (Lockwood & Tsybeskov, 2004).


3.4.1 Disordered Si QDs in
insulating matrix


The
transport properties of the ensembles of disordered Si QDs in insulating matrix
could be explained in terms of the percolation theory, which has already been
successfully implemented to explain the transport processes in granular metals
(Abeles et al., 1975).









	
  

 	
  

 
	
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Indeed,
this theory describes the effect of the system’s connectivity on its
geometrical and physical properties. In the case of granular metals, in a
system of N metallic spheres
embedded randomly in an insulating matrix, there will be a critical density of
spheres Nc above which
a “continuous” metallic network will be formed and a metallic bulk-like
conduction will dominate. Correspondingly, Nc
is the classical percolation threshold (Fonseca & Balberg, 1993)
(Balberg et al., 2004). For N < Nc,
the electron transfer between the individual grains is possible only by
tunneling (Abeles et al., 1975) (Balberg et al., 2004).


To
what concerns the ensemble of Si QDs, there can be distinguished five different
structural-electrical regimes, such that in each of them we may expect a
different transport mechanism to dominate. These regimes are (a) uniformly dispersed in insulating
matrix isolated spherical QDs; (b) the
transition regime, where some of the QDs starts to “touch” their neighbors; (c) the intermediate regime, where clusters
of “touching” QDs are formed; (d) the
percolation transition regime where the above clusters form a global continuous
network; and (e) the regime where
the percolation cluster of “touching” QDs is well formed and geometrically
non-”touching” QDs are rarely found. Fig. 15 (a), (b) and (c) present typical
examples of ensembles of Si QDs corresponding to regimes (a), (c)
and (e), respectively.
It is worth noting that the connectivity between “touching” QDs in ensembles of
semiconductor QDs is different than in granular metals. Usually there are
narrow (no more than 0.5 nm wide) boundaries formed between the nanoparticles,
which involves at least a different crystallographic orientation of the
touching crystallites. This quantum size “grain boundary” limit has not been
studied so far. In a literature the charge transfer process between such
“touching” QDs was termed as “migration” (Antonova et al., 2008) (Balberg et
al., 2007).





Fig.
15. HRTEM images of the ensembles of Si QDs corresponding to different
structural-electrical regimes: (a) uniformly dispersed isolated spherical QDs
(regime a), (b) clusters of
“touching” QDs (regime c) and (c)
percolation clusters of “touching” QDs (regime e) (Antonova et al., 2008).


The
effect of the connectivity on the transport properties (dark and
photoconductivity) of the ensembles of Si QDS is illustrated on Fig.
16. As one can see, the global picture of transport in Si QDs ensembles is
reminiscent of that of granular metals, but the details are quite different.
For the samples with low Si content (related to the number of Si QDs in the
ensemble), which are characterized by the isolated Si QDs (regime a), the local conductivity









	
  

 	
  

 
	
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is
determined by the tunneling of charge carriers under Coulomb blockage2
between adjacent nanocrystallites similar to the case encountered in granular
metals in the dielectric regime (Abeles et al., 1975) (Balberg et al., 2004).
Indeed, as long as Si QDs or clusters of Si QDs are small enough, they “keep”
the carrier that resides in them and become charged when an excess charge
carrier reaches them. Hence, the transport through the system can take place
only if a corresponding charging (or Coulomb) energy is provided.





Fig.
16. Dependence of the dark conductivity and the photoconductivity on the Si
content (related to the number of Si QDs in ensemble)


With
increasing of Si content (Fig. 15 (b)), the interparticle distance decreases
and the tunneling-connected quantum dot clusters grow in size. The
“delocalization” of the carrier from its confinement in the individual quantum
dot to larger regions of the ensemble will take place, i.e., the charge carrier
will belong to a cluster of QDs rather than to an individual QD.
Correspondingly, this will also yield a decrease in the local charging energy
in comparison with that of the isolated QD and the distance to which the charge
carrier could wander will increase and as a consequence the conductivity of the
ensemble will increase as well. The charge carrier transport in the case of
regime (c) is thus determined by
the intracluster migration and by the intercluster tunneling.


As
one can see from the Fig. 16, the maximal possible conductivity is assured in
the case of highly percolating system of Si QDs (regime (e)). However, the conduction in this
regime is quite different from that one of the granular metals since there are
still boundaries between touching Si QDs. In fact, the corresponding migration
process is similar to that in polycrystalline semiconductors, but now the
boundaries are on the quantum scale. It was suggested that in this regime the
migration dominates the transport properties and the global conductivity is
limited by the interface between the touching QDs (Balberg, 2010).


	
  

 
	
 _______________

 
	
 2 The transfer of an electron from a given
 neutral particle to an adjacent neutral particle, charges this particle by
 one (positive) elementary charge (q) and
 that of the adjacent particle by one (negative) elementary charge. If the
 capacitance of the individual particle in its corresponding environment is C0, the energy needed to be
 supplied for the above “electron-hole” transfer by tunneling is then: E=q2/C0. This
 energy, which opposes to the transfer of charge carriers, is known as the
 Coulomb blockage energy, which is of the order of a tenth of an eV. In
 general, one can say, that a tunneling process is thermally activated when it
 requires a supply of energy.

 













	
  

 	
  

 
	
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Form
the photovoltaic point of view, the thin films constituted of interconnected Si
QDs are the most promising candidates for higher bandgap materials in all-Si
tandem cell approach. Indeed, the highly percolating system of Si QDs will
ensure the most favorable conditions for the electronic transport between the
nanocrystals and, as it was discussed previously, the bandgap value in such
structures could be adjusted in the large range covering the major part of the
solar spectrum (Nychyporuk et al., 2009).


3.4.2 Nanocrystalline Si
superlattices


Nanocrystalline
Si superlattices have been proposed as candidates for the high bandgap absorber
component in all-Si tandem solar cells. They consists of thin dielectric and Si
QD based layers alternating in one direction, i.e., heterostructure type –I
superlattices. The period of such a superlattice usually is much larger than
the lattice constant but is smaller than the electron mean free path. Such a
structure possesses, in addition to a periodic potential of the crystalline
lattice, a potential due to the alternating semiconductor layers. The existence
of such a potential significantly changes the energy bandstructure of the
semiconductors from which the superlattice is formed. The coupling among QDs
occurs, leading to a splitting of the quantized carrier energy levels of single
dots and formation of three-dimensional minibands (Lazarenkova & Balandin,
2001) (Jiang & Green, 2006) as sketched on Fig. 17 (shaded areas).


The
charge carrier mobility, which has a crucial impact on a charge-collection
efficiency in solar cells, depends on the dominant transport regime at given
operating conditions, which may be described by mini-band transport, sequential
tunneling or Wannier-Stark hopping (Wacker, 2002). The sequential resonant
tunneling (SRT) was suggested to be the most prominent for efficient carrier
collection in Si QD solar cells (Raisky et al., 1999).


The
schematic view of the sequential resonant tunneling transfer in a
multiple-quantum-well structure is depicted on Fig. 18. Electrons tunnel from
the ground state of the jth
well into an excited state of the (j+1)thwell. The tunneling process is then followed by intrasubband energy
relaxation from the excited state to the ground state. This two-step scheme can
be repeated as many times as needed to build the required thickness for optimal
solar absorption. Resonance occurs when the E2-E1
=│qFd│ condition is satisfied where E1 and E2 are the ground and first
excited subband energies of the quantum well, q
the electron charge, F the
internal electric field, and d the
spatial structural period. From Fig. 18 it is clear that the bottleneck of
electron transfer is the last (Nth)
well, where carriers have to transfer through a significantly thicker right
barrier than inside the multiple-quantum-well region. This can lead to charge
build up and, consequently, screening of the built-in field. The
photogeneration and transport in superlattice absorbers, on the example of a
Si-SiOx multilayer structure embedded in the intrinsic region of a
p-i-n diode was recently numerically investigated (Aeberhard, 2011). The model
system under investigation is shown on the Fig. 19. It consists of a set of
four coupled quantum wells of 6 monolayer3 (ML) width with layers
separated by oxide barriers of 3-ML thickness, embedded in the intrinsic region
of a Si p-i-n diode. The doping density was 1018 cm-3 for
both electrons and holes. Insertion of the oxide barriers leads to an increase
of the effective bandgap in the central region of the diode from 1.1 to 1.3 eV.
The spectral rate of carrier generation in the confined states under
illumination with monochromatic light at photon energy 1.65 eV and intensity of
10 kW/m2


	
  

 
	
 _______________

 
	
 3 The
 monolayer thickness is half the Si lattice constant, i.e., 2.716 A

 













	
  

 	
  

 
	
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is
shown on the Fig. 20. At this photon energy both the lowest and the second
minibands are populated. The photocurrent originating in this excitation is
shown on Fig. 21. Current flows also in first and second minibands, which means
that relaxation due to scattering is not fast enough to confine transport to
the band edge. However, transport of photocarriers is strongly affected by the
inelastic interactions, and is the closest to the sequential tunneling regime.
We can thus conclude that in the case of high internal fields, excess charge is
transported via sequential tunneling in the miniband where it is generated.


The
sequential resonance tunneling enhances the photocarrier collection and reduces
radiative recombination losses (Raisky et al., 1999). However it should be
noted that SRT increases both photocurrent (Iph) and dark current (Idc). The
total current of a photovoltaic device is the difference of these two currents,
and thus, to take advantage of SRT, a solar cell possessing a superlattice
structure should be designed to have the resonance in the region where Iph" Idc.


The
main challenge of the tandem structure is to achieve sufficient carrier
mobility and hence a reasonable conductivity. This generally requires formation
of a true superlattice with overlap of the wave function for adjacent QDs,
which in turns requires either close spacing between Si nanocrystals or low
barrier height. Transport properties strongly depend on the matrix in which the
Si quantum dots are embedded. Indeed, the electron or the hole wavefunctions
exponentially decay with distance. Fig. 22 shows the penetration length of the
wave function of electron of a single quantum well into different high-bandgap
materials having different barrier heights. As one can see, the tunneling
probability heavily depends on the barrier height. The penetration length is
bigger for the materials with lower barrier height. Thus Si3N4
and SiC giving lower barriers than SiO2, allow larger dot spacing for
a given tunneling current. For example, the QDs in SiO2 matrix would
have to be separated by no more than 1-2 nm of matrix, while they could be
separated by more than 4 nm of SiC.


The
influence of the fluctuations in spacing and size of the QDs on the carrier
mobility was also investigated (Jiang & Green, 2006). The calculations have
shown that the interdot distance has larger impact on the calculated carrier
mobility while the dot size can be used to control the band energy level.


3.5 Fabrication of Si QD PV
devices


Recently
it have been reported on the realization of interdigitated silicon QD solar
cells on quartz substrate(Conibeer et al., 2010). Schematic view of the
fabricated solar cells is shown on the Fig. 23 (a). The p-n diodes were
fabricated by sputtering alternating layers of SiO2 and SRO onto
quartz substrate with in situ B and P doping. The top B doped bi-layers were
selectively etched to create isolated p-type mesas and to access the buried P
doped bi-layers. Aluminium contacts were deposited by evaporation, patterned
and sintered to create ohmic contacts on both p- and n-type layers. The area of
the fabricated interdigitated solar cells was 0.12 cm-2. One of the
derivatives of the presented approach was the fabrication of the p-i-n diodes.
Indeed, it is expected that the intrinsic layer will have a longer lifetime
that the doped material leading to an improved photocurrent.


I-V
measurements in the dark and under 1-sun illumination (Fig. 23 (b)) indicate a
good rectifying junction and generation of an open-circuit voltage, VOC,
up to 492 mV (Conibeer, 2010). The high sheet resistance of the deposited
layers, in conjunction with the insulating quartz substrates, causes an
unavoidable high series resistance in the device.









	
  

 	
  

 
	
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Fig.
17. The energy bandstructure of a semiconductor type-I heterostructure
superlattice : Eg1 and Eg2 are the bandgaps, Ec1
and Ec2 are the bottoms of the conduction bands of narrow bandgap
and wide bandgap semiconductors, respectively; d is the period of the
heterostructure superlattice (Mitin, 2010).





Fig.
18. Sequential resonant tunneling transfer in multiple-quantum-well structure. E1 and E2 are the energies of the
ground and first excited states in the quantum well, respectively, and d the superlattice period (Raisky et al.,
1999).





Fig. 19. Spatial structure and doping profile of the p-i-n model system
(Aeberhard, 2011).





Fig.
20. Spatially and energy resolved charge carrier photogeneration rate in the
quantum well region at short-circuit conditions and under monochromatic
illumination with energy of 1.65 eV and intensity of 10kW/m2
(Aeberhard, 2011).









	
  

 	
  

 
	
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Fig.
21. Spatially and energy-resolved charge carrier short-circuit photocurrent
density in the quantum well region under monochromatic illumination with energy
of 1.65 eV and intensity of 10kW/m2 (Aeberhard, 2011).





Fig.
22. Penetration length of the wave function of confined electron into barrier
layers. Vb represents a barrier height for each barrier material
(Aeberhard, 2011).





Fig.
23. (a) Schematic representation of the fabricated interdigitated devices
(Conibeer et al., 2010); (b) Dark and illuminated I-V measurements of p-i-n
diodes with 4 nm SRO/2 nm SiO2 bilayers.









	
  

 	
  

 
	
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The
high resistance severely limits both the short-circuit current and the fill
factor of the cells, particularly under illumination. Significant improvement
is expected once the parasitic series resistance is eliminated.


Further
evidence that this photovoltaic effect occurs in a material with an increased
bandgap is given by temperature dependent I-V mesurements, from which an
electronic bandgap for the Si QD nanostructure materials can be extracted. A
bandgap of 1.8 eV was extracted for a structure containing Si QDs with a
nominal diameter of 4 nm. However, this value will be due to a combination of
other components in series with the material bandgap, hence the true material
bandgap will be less than 1.8 eV.


Homojunction
Si QD devices have also been fabricated in a SiC matrix using the superlattice
approach (Song et al., 2009). Fig. 24 (a) shows a schematic diagram of a n-type
Si QD: SiC/ p-type Si QD:SiC homojunction solar cell fabricated on a quartz
substrate. The n-type Si QD emitter was approximately 200 nm thick and the
p-type base layer is approximately 300 nm thick. This devices have given VOC
of 82 mV that is promising initial value for a Si QDs in SiC solar cells on
quartz substrate. Improvement of the device structure and optimization of
dopant incorporation is expected to improve this value.





Fig.
24. (a) Schematic diagram of a n-type Si QD: SiC/ p-type Si QD:SiC/ quartz
homojunction solar cell (Song et al., 2009); (b) The concept of the transport
improvement: alternating layers of Si3N4 and SRO.


Current
in both these SiO2 and SiC matrix devices was very small, due
principally to the very high lateral resistance and also because of the small
amount of absorption in the approximately 200 nm of material used. Indeed, as
it was discussed in the previous paragraph, transport in these devices relies
on tunneling and hopping between adjacent QDs. To maximize the tunneling
probability, the barrier heights between QDs must be low, but this then
compromises the degree of quantum confinement and the height of the confined
energy levels and hence the effective bandgap obtained. The solution is to
introduce anisotropy between the growth, z,
direction and the x-y plane. This
can achieved by maintaining strong confinement through the use of a large
barrier height oxide in the plane, but to intersperse these layers with layers
of lower barrier height such as Si3N4 or SiC, thus giving
higher tunneling probability in the z direction (Fig. 24 (b)) (Di et al.,
2010). The very first results on this approach were rather promising and showed
the decreasing in the vertical resistivity of such Si QD nanostructures with
SiNx interlayers (Di et al., 2010). To what concerns the increasing
of the VOC the most potential route is the passivation of the
defects throug+h the hydrogenation.









	
  

 	
  

 
	
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The
first prototypes of Si QD PV devices were successfully developed. Up today they
present VOC, ISC and fill factor (FF) values which still
lower than those ones of the 1st generation PV cells based on bulk
Si – but all these problems are being addressed. The next step implies the
further optimization of the fabrication parameters, developing of the efficient
doping technique and defect passivation.


4. Silicon nanowire solar cells


Nanowire
solar cells demonstrated to date have been primarily based on hybrid
organic-inorganic materials or have utilized compound semiconductors such as
CdSe. Huynh et al. utilized CdSe
nanorods as the electron-conducting layer of a hole conducting polymer-matrix
solar cell (Huynh et al., 2002) and produced an efficiency of 1.7 % for AM 1.5
irradiation. Similar structures have been demonstrated for dye-sensitized solar
cells using titania or ZnO nanowires, with efficiencies ranging from 0.5 % to
1.5 % (Law et al., 2005). These results show the benefits of using nanowires
for enhanced charge transport in nanostructured solar cells compared to other
nanostructured architectures. The Si nanowires (Si NW) solar cells have a
potential to provide the equal or better performance to crystalline Si solar
cells with processing methods similar to thin film solar cells (Tsakalakos et
al., 2007) (Uchiyama et al., 2010) (Andra et al., 2008).


4.1 Fabrication of Si nanowires


The
techniques to produce nanowires are normally divided into (i) bottom up and (ii) top down approaches.


4.1.1 Bottom up approach


The
bottom up approach starts with individual atoms and molecules and builds up the
desired nanostructures. One of the mostly used methods in this family of
approaches is the Vapor Liquid Solid (VLS) method which uses metal
nanotemplates on Si wafer or on Si thin film (Kelzenberg et al., 2008) (Tian et
al., 2007) (Tsakalakos et al., 2007, 2007a). In this method a liquid metal
cluster acts as the energetically favored site for vapor-phase reactant
absorption and when supersaturated, the nucleation site for crystallization. An
important feature of this approach for nanowire growth is that phase diagrams can
be used to choose a catalyst material that forms a liquid alloy with the
nanowire material of interest, i.e. Si in our case. Also, a range of potential
growth temperatures can be defined from the phase diagram such that there is
coexistence of liquid alloy with solid nanowire phase. The main advantages of
the VLS method which should be cited are the rather high growth rate of several
100 nm/min and the fact that perfectly single crystalline nanowires form.


A
schematic diagram illustrating the growth of Si nanowires by the VLS mechanism
is shown on (Fig. 25 (a)). When the nanocatalysts become supersaturated with
Si, a nucleation event occurs producing a solid/liquid Si/Au-Si alloy
interface. In order to minimize the interfacial free energy, subsequent solid
growth/crystallization occurs at this initial interface, which thus imposes the
highly anisotropic growth constrain required for producing nanowires. Its
growth continues in the presence of reactant as long as catalyst nanocatalyst
remains in the liquid state. Typically the growth is performed by using SiH4
as the Si reactant, and diborane and phosphine as p- and n-type dopants,
respectively. The growth can be carried out using Ar, He or H2 as
carrier gas, which enables an added degree of freedom for the nanowire growth.
For exemple, the use of H2 as a carrier gas can









	
  

 	
  

 
	
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passivate
the growing solid surface and reduce the roughness (Wu et al., 2004) while Ar
and He can enhance radial deposition of a specific composition shell.


Most
frequently gold is taken as a template for nanowire growth. The Au-Si binary
phase diagram (Fig. 25 (b)) predicts that Au nanocatalysts will form liquid
alloy droplets with Si at temperatures higher than the eutectic point which is
3630C. The Au nanocatalysts can be prepared either from commercially available
gold colloids or by depositing a thin Au film followed by a heating step above
the eutectic temperature during which a Au-Si liquid film forms to disintegrate
into nano-droplets. Another simple and effective method for producing metal
nanoparticles (Au, Ag...) at room temperature is based on their electroless
deposition on the surface of Si or hydrogenated Si nitride films (Nychyporuk et
al., 2010). It should be mentioned that it is still a controversial issue how
gold is incorporated into the wire and thus how it influences the electronic
properties of the nanowire. Gold is a deep-level defect in bulk Si and if it is
also true for nanowires grown from Au droplets. Hence the alternative metals
like In, Sn, Al (Ball et al., 2010) are actually under investigation for using
as nanocatalysts during the nanowire growth.


Si
NWs grow with a diameter similar to that of the template droplet. The nanowire
diameters are on average 1-2 nm larger than the starting nanocatalyst size. As
a result a carpet of perfect single crystalline NWs of 10 to 200 nm in diameter
and several micrometers in length can be grown on the crystalline substrate
(Fig. 26 (a)) (Andra 2008). High-resolution transmission electron microscopy
(HRTEM) was used to define in detail the structures of these nanowires (Wu et
al., 2004). As synthesized Si NWs are single crystalline nanostructures with
uniform diameters. Studies of the ends of the nanowires show that they often
terminate with Au nanoparticles (Fig. 26 (b)). In addition, the
crystallographic growth directions of Si NWs have also been investigated using
HRTEM and systematic measurements reveal that the growth axes are related to
their diameters (Wu et al., 2004) (Cui et al., 2001a). For diameters between 3
and 10 nm, 95% of the Si NWs were found to grow along the <110>
direction, for diameters between 10 and 20 nm, 61% of the Si NWs grow along the
<112> direction, and for diameters between 20 and 30 nm, 64% of the Si
NWs grow along the <111> direction. These results demonstrate a clear
preference for growth along the <110> direction in the smallest Si NWs
and along <111> direction in larger ones (Zhong et al., 2007).
Cross-sectional HRTEM analysis has revealed that the nanowires could have
triangular, rectangular and hexagonal cross section with well – developed
facets (Vo et al., 2006) (Jie et al., 2006) (Zhang et al., 2005).


The
VLS method permits to fabricate the nanowires of well-defined length with
diameters as small as 3 nm (Wu et al., 2004). The electronic properties can be
precisely controlled by introducing dopant reactants during the growth.
Addition of different ratios of diborane or phosphine to silane reactant during
growth produces p- or n- type Si nanowires with effective doping concentrations
directly related to the silane: dopant gas ratios (Cui et al., 2000) (Cui et
al., 2001) (Zheng et al., 2004) (Fukata, 2009). It was demonstrated that B and
P can be used to change the conductivity of Si NWS over many orders of
magnitude (Cui et al., 2000). The carrier mobility in SiNWs can reach that one
in bulk Si at a doping concentration of 1020 cm-3 and
decreases for smaller diameter wires. Temperature dependent measurements made
on heavily doped SiNWs show no evidence for single electron charging at
temperatures down to 4.2 K, and thus suggest that SiNWs possess a high degree
of structural and doping uniformity. Moreover, TEM studies of boron- and
phosphorous doped SiNWs have shown that contrary to Si QDs the introduction of
impurity atoms during the nanowire growth does not change their crystallinity.
The ability to prepare well-defined doped nanowire during synthesis
distinguishes nanowires from QDs.









	
  

 	
  

 
	
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Fig.
25. (a) Schematic diagram illustrating the growth of Si nanowires by the VLS
mechanism. (b) Binary phase diagram of Au-Si (Zhong et al., 2007).





Fig.
26. (a) SEM image of a Si NW carpet grown from Au nanoparticles (150 nm
diameter) on a multicrystalline Si wafer (Andra 2008); (b) HRTEM image of the
gold catalyst/nanowire interface of a Si NW with a <110> growth axis.
Scale bar is 5 nm (Wu et al., 2004).


4.1.2 Top down approach


In
top down approach relies on dimentional reduction through selective etching and
various nanoimprinting techniques. For example, well aligned Si NW arrays can
be obtained by electroless metal-assisted chemical etching in HF/AgNO3
solution. Basically, a noble metal is deposited on the surface in the form of
nanoparticles which act as catalyst for Si etching in HF solution containing an
oxidizing agent. As a consequence, the etching only occurs in the vicinity of
the metal nanoparticles and results in the formation of well defined mesopores
(20-100 nm in diameter (Fang et al., 2008) (Peng et al., 2005) (Fig. 27 (a)).


One
of the main advantages of the top – down methods as compared to the bottom –up
is that it is possible to start the processing by using conventional wafers
with already performed diffusion regions (ex. p+nn+ ) and
then etch it into SiNWs. Each nanowire will









	
  

 	
  

 
	
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thus
present the p-n junction already formed. However this method presents some
drawbacks. One of them is rather poor size control.


To
overcome this problem colloidal crystal patterning combined with metal-assisted
etching was applied (Wang et al., 2010) (Wang et al., 2010). The main idea is
the next one. The sphere lithography is based on the self-organization of
micrometer/nanometer spheres into a monolayer with a hexagonal close-packed
structure. Typical material used for the spheres are silica and polystyrene,
which are commercially available with narrow size distribution. The deposition
of a monolayer of the spheres on a substrate is used as a patterning mask for
thin metallic film evaporated on the Si wafer. After the sphere dissolution,
the Si surface that comes in contact with the metal is selectively etched,
leaving behind an array of Si NWs whose diameter is predefined by the size of
holes in the metal film, while the length is determined by the etching time.
This method enables the formation of large scale arrays with long – range
periodicity of vertically standing nanorods/nanowires with well controlled
diameter, length and density (Fig. 27 (b)). It should be however noted that
this methods does not permit to achieve SiNWs with the diameters lower than 50
nm and thus potentially cannot be applied for all-Si tandem solar cell.





Fig.
27. SEM images of Si NWs obtained by (a) simple metal-assisted etching
technique (Andra 2008) and (b) by colloidal crystal patterning combined with
metal-assisted etching approach (Wang et al., 2010).


4.2 Optical properties of Si
nanowires


A
variety of the optical techniques have shown that the properties of nanowires
are different to those of their bulk counterparts, however the interpretation
of these measurements is not always straightforward. The wavelength of light
used to probe the sample is usually smaller than the wire length, but larger
than the wire diameter. Hence, the probe light used in the optical measurement
cannot be focused solely onto the nanowire and the wire and the substrate on
which the wire rests (or host material if the wire is embedded in a template)
are probed simultaneously. For example, for measurements such as
photoluminescence, if the substrate does not luminesce or absorb in the
frequency range of the measurements, PL measures the luminescence of the
nanowire directly and the substrate can be ignored. In reflection and
transmission measurements, even a non-absorbing substrate can modify the
measured spectra of nanowires. However, despite these technical difficulties it
was









	
  

 	
  

 
	
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experimentally
proved that Si nanowire materials have exhibited properties such as ultrahigh
surface area ratio, low reflection, absorption of wideband light, and tunable
bandgap.


4.2.1 Bandgap of quantum SiNWs


In
2003, scanning tunneling spectroscopy measurements on individual oxide-removed
Si NWs showed that the optical gap of Si NWs increased with decreasing of Si NW
diameter from 1.1 eV for 7 nm to 3.5 eV to 1.3 nm (Ma et al., 2003). Since, a
large number of theoretical and experimental works have been done to explore
the effect of the chemical passivation, surface reconstruction, cross section
geometry and growth orientation on electronic structure of SiNWs
(FernÁndez-Serra et al., 2006) (Yan et al., 2007) (Vo et al, 2006). For
example, it was shown than the bandgap of [110] SiNWs is the smallest among those
of the [100], [112], and [111] wires of the same diameter (Cui et al, 2000). It
should be also noted that the magnitude of the energy increase/decrease in
SiNWs induced by quantum confinement is different for each point of the
bandstructure. It was predicted that the conduction-band-minimum (CBM) energy
increases more near the X point than near the I. Therefore for nanowires with
sufficiently small dimensions, this difference in the energy shifts at
different points in the Brillouin zone is large enough to move the CBM at the X
point above the CBM at the Y point
(Vo et al., 2006). Then a transition from an indirect to direct gap material
occurs. The indirect to direct transition does not depend on the special
cross-sectional shapes and the bandgaps of [110] SiNWs remain direct event for
SiNWs with dimensions up to 7 nm. The dependence of the bandgap on SiNW
dimension D is shown on the Fig.
28 (a). It is obvious that the smaller the dimension of the nanowire the larger
the bandgap due to quantum confinement. As D
decreases from 7 to 1 nm, the bandgap increases from 1.5 to 2.7 eV.
In addition to the size dependence, the energy gap also shows significantly
different change with respect to the cross-sectional shape. The bandgaps of
rectangular and hexagonal SiNWs are rather close while distinctly smaller than
that of the triangular SiNWs.





Fig. 28. (a) Bandgap of SiNWs versus the transverse dimension D. (b) Bandgap of SiNWs versus SVR (Yao et
al., 2008).


The
significant cross-sectional shape effects on band gap and size dependence can
be understood from the concept of surface –to-volume ratio (SVR). Because of
the quantum confinement effect, the band gap increases as the material
dimension is reduced, thus









	
  

 	
  

 
	
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leading
to an increase of SVR. In other words, SVR has the impact of enlarging band
gap. At the same transverse dimension, triangular SiNW has larger SVR than
those of the rectangular SiNW and hexagonal SiNW As a result, its larger SVR
induces the largest band gap among those of the rectangular and hexagonal SiNWs
and the strongest size dependence. The bandgap values versus SVR of the SiNWs
are shown in Fig. 28 (b). The SVR effect on the bandgap of [110] SiNWs with any
cross-sectional shape and area can be described by a universal expression (Yao
et al, 2008)


EG (eV)=1.28+0.37 x S (nm-1),


where
S is the value of the SVR in unit
of nm-1. The bandgap of SiNWs are usually difficult to measure, but
their transverse cross-sectional shape and dimension are easy to know, so it is
of significance to predict the bandgap values of SiNWs by using the above
expression.


4.2.2 Optical reflection and
absorption in SiNWs


Si
NW PV devices show improved optical characteristics compared to planar devices.
Fig. 29 (a) shows typical optical reflectance spectra of SiNW film as compared
to solid Si film of the same thickness (10mm) (Tsakalakos et al., 2007a). As
one can see, the reflectance of the nanowire film is less than 5% over the
majority of the spectrum from the UV to the near IR and begins to increase at
7000nm to a values of 41% at the Si band edge (1100 nm), similar to the bulk
Si. It is clear that the nanowires impart a significant reduction of the
reflectance compared to the solid film More striking is the fact that the
transmission of the nanowire samples is also significantly reduced for
wavelength greater than 7000nm (Fig. 29 (b)). This residual absorption is
attributed to strong IR light trapping4 coupled with the presence of the
surface states on the nanowires that absorb below bandgap light. However, the
level of optical absorption does not change with passivation, which further
indicates that light trapping plays a dominant role in the enhanced absorption
of the structures at all wavelength. It should be also noted that the
absorption edge of a nanowire film shifts to longer wavelength and approaches
the bulk value as the nanowire density is increased. Essentially, the Si
nanowire arrays act as sub-wavelength cylindrical scattering elements, with the
mactroscopic optical properties being dependent on nanowire pitch, length, and
diameter.





Fig.
29. Total (a) reflectance and (b) transmission data from integrated sphere
measurements for 11 mm thick solid Si film and nanowire film on glass substrate
(Tsakalakos et al, 2007).


	
  

 
	
 _______________

 
	
 4 Light trapping is typically defined as the
 ratio of the effective path length for light rays confined within a structure
 with respect to its thickness.

 













	
  

 	
  

 
	
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4.3 Electrical
transport in SiNWs


Important
factors that determine the transport properties of Si nanowires include the
wire diameter (important for both classical and quantum size effect), surface
conditions, crystal quality, and the crystallographic orientation along the
wire axis (Ramayya et al., 2006) (Duan et al., 2002).


Electronic
transport phenomena in Si nanowires can be roughly divided into two categories:
ballistic transport and diffusive transport. Ballistic transport phenomena
occur when the electrons can travel across the nanowire without any scattering.
In this case the conduction is mainly determined by the contact between the
nanowire and the external circuit. Ballistic transport phenomena are usually
observed in very short quantum wires. On the other hand, for nanowires with
length much larger than the carrier mean free path, the electrons (or holes)
undergo numerous scattering events when they travel along the wire. In this
case, the transport is in the diffusive regime, and the conduction is dominated
by carrier scattering within the wires, due to lattice vibrations, boundary
scattering, lattice and other structural defects and impurity atoms.


The
electronic transport behavior of Si nanowires may be categorized based on the
relative magnitudes of three length scales: carriers mean free path, the de
Broglie wavelength of electrons, and the wire diameter. For wire diameters much
larger than the carrier mean free path, the nanowiers exhibit transport
properties similar to bulk materials, which are independent of the wire
diameter, since the scattering due to the wire boundary is negligible, compared
to other scattering mechanisms. For wire diameters comparable or smaller than
the carrier mean free path, but still larger than the de Broglie wavelength of the
electrons, the transport in the nanowire is in the classical finite regime,
where the band structure of the nanowire is still similar to that of the bulk,
while the scattering events at the wire boundary alter their transport
behavior. For wire diameters comparable to electronic wavelength (de Broglie
wavelength of electrons), the electronic density of states is altered
dramatically and quantum sub-bands are formed due to quantum confinement effect
at the wire boundary. In this regime, the transport properties are further
influenced by the change in the band structure. Therefore, transport properties
for nanowires in the classical finite size and quantum size regimes are highly
diameter-dependent. Experimentally it was shown that the carrier mobility in SiNWs
can reach that one in bulk Si at a doping concentration of 1020 cm-3
and decreases for smaller diameter wires (Cui et al., 2000).


Because
of the enhanced surface-to-volume ratio of the nanowires, their transport
behavior may be modified by changing their surface conditions. For example, it
was shown on the n-InP nanowires, that coating of the surface of these
nanowires with a layer of redox molecules, the conductance may be changed by
orders of magnitude (Duan et al., 2002).


4.4 Comparison
of axial and radial p-n junction nanowire solar cells


Independently
of the nanowire preparation method two designs of NW solar cells are now under
consideration with p-n junction either radial or axial (Fig. 30). In the radial
case the p-n junction covers the whole outer cylindrical surface of the NWs.
This was achieved either by gas doping or by CVD deposition of a shell
oppositely doped to the wire (Fang, 2008) (Peng, 2005) (Tian 2007). In the
axial variant, the p-n junction cuts the NW in two cylindrical parts and require
minimal processing steps (Andra 2008). However, solar cells that absorb photons
and collect charges along orthogonal directions meet the optimal relation
between the absorption values and minority charge carrier diffusion lengths
(Fig. 30 (a)) (Hochbaum 2010). A solar cell consisting of arrays of radial p-n
junction nanowires (Fig. 30 (b)) may









	
  

 	
  

 
	
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provide
a solution to this device design and optimization issue. A nanowire with a p-n
junction in the radial direction would enable a decoupling of the requirements
for light absorption and carrier extraction into orthogonal spatial directions.
Each individual p-n junction nanowire in the cell could be long in the
direction of incident light, allowing for optimal light absorption, but thin in
another dimension, thereby allowing for effective carrier collection.





Fig.
30. Schematic views of the (a) axial and (b) radial nanowire solar cell. Light
penetration into the cell is characterized by the optical thickness of the
material (a is the absorption coefficient), while the mean free path of
generated minority carriers is given by their diffusion length. In the case of
axial nanowire solar cell, light penetrates deep into the cell, but the
electron-diffusion length is too short to allow the collection of all
light-generated carriers (Kayes et al., 2005).


The
comparison between the axial and radial p-n junction technologies for solar
cell applications was performed in details in Ref (Kayes et al, 2005). In the
case of radial p-n junction, the short-circuit current (Isc) increases with the
nanowire length and plateaus when the length of the nanowire become much
greater than the optical thickness of the material. Also, Isc was essentially independent
on the nanowire radius, provided that the radius (R) was less than the minority carrier diffusion length (Ln). However, it decreases
steeply when R > Ln. Isc
is essentially independent of trap density in the depletion region.
Being rather sensitive to a number of traps in the depletion region, the open
circuit voltage Voc decreases
with increasing nanowire length, and increases with nanowire radius. On the
other hand the trap density in the quasineutral regions had relatively less
effect on Voc. The
optimal nanowire dimensions are obtained when the nanowire has a radius
approximately equal to Ln and
a length that is determined by the specific tradeoff between the increase in Isc and the decrease in Voc with length. In the case of
low trap density in the depletion region, the maximum efficiency is obtained
for nanowires having a length approximately equal to the optical thickness. For
higher trap densities smaller nanowire lengths are optimal.


Radial
p-n junction nanowire cells trend to favor high doping levels to produce high
cell efficiencies. High doping will lead to decreased charge-carrier mobility
and a decreased depletion region width, but in turn high doping advantageously
increases the build-in voltage. Because carriers can travel approximately one
diffusion length through a quasineutral region before recombining, making the
nanowire radius approximately equal









	
  

 	
  

 
	
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to
the minority -electron diffusion length allows carriers to traverse the cell
even if the diffusion length is low, provided that the trap density is
relatively low in the depletion region.


An
optimally designed radial p-n junction nanowire cell should be doped as high as
possible in both n- and p- type regions, have a narrow emitter width, have a
radius approximately equal to the diffusion length of the electrons in the
p-type core, and have a length approximately equal to the thickness of the
material. It is crucial that the trap density near the p-n junction is
relatively low. Therefore one would prefer to use doping mechanisms that will
getter impurities away from the junction. By exploiting the radial p-n junction
nanowire geometry, extremely large efficiency gains up to 11% are possible to
be obtained.


4.5 Fabrication of Si QD PV
devices


By
using VLS method (Tian et al, 2007) (Kelzenberg et al, 2008) (Rout et al, 2008)
(Fang et al., 2008) (Perraud et al., 2009) as well as by the etching method
(Garnett et al., 2008) (Peng et al., 2005). SiNW based photovoltaic devices
were experimentally demonstrated. Nearly all the works were concerned with Si
wafers as a substrate. However, it should be noted that for competitive solar
cells, low cost substrates, such as glass or metal foils are to be preferred.
Schematic view of the VLS fabricated structure of the SiNW array solar cells is
illustrated on Fig. 31 (a). The n-type SiNWs were prepared by the VLS method on
(100) p-type Si substrate (14-22 Wcm). Device fabrication started from the
evaporation of 2-nm thick gold film followed by annealing at 5500C for 10 min
under H2 flow to form Au nanocatalyzers. SiNWs were subsequently
grown at 5000 with SiH4 diluted in H2 as the gas
precursor. N-type doping was achieved by adding PH3 to SiH4,
with PH3/SiH4 ratio of 2x10-3 corresponding to
a nominal phosphorous density of 1020cm-3. After the VLS
growth the gold catalysts were etched off in KI/I2 solution, and the
doping impurities were activated by rapid thermal annealing at 7500 for 5 min.
The SiNW array was then embedded into spin-on-glass (SOG) matrix. Indeed, SOG
matrix ensures a good mechanical stability of the SiNW array and enables
further processing steps, such as front surface planarization and electrical
contact deposition. The planarization step is normally performed by the
chemical-mechanical polishing. To form the front contacts indium-tin-oxide
(ITO) was firstly deposited on planarized SOG surface followed by the
deposition of Ni/Al contact grid. As back electrical contact, the sputtered and
annealed Al was used. The area of the fabricated SiNW solar cell was 2.3 cm-2.


The
sheet resistance of n-type SiNWs embedded into SOG matrix was estimated to be
10-4 W/sq. I-V measurements in the dark and under 1-sun illumination
(Fig. 31 (b)) indicate a good rectifying junction. The measured ISC, VOC
and FF were 17 mA/cm2, 250mV and 44%, respectively, leading to an
energy-conversion efficiency of 1.9%. The VOC of Si NW solar cell
was shown to be increased up to 580 mV (Peng et al, 2005). The parasitic series
resistance found for SiNW solar cells (5 W cm-2) was slightly
larger than in the standard 1st generation solar cells (2 W cm-2),
however the p-n junction reverse current was of the order of 1 mA/cm2
with is about 100 times bigger than in typical Si solar cells (1 pA/cm2).
Such a high pn junction reverse current indicates a high density of localized
electronic states within the bandgap, which act as generation-recombination
centers. These states may come from contamination of Si by gold which is used
as catalyst for VLS growth. Other types of metallic catalysers, like Sn, were
also used (Uchiyama et al, 2010).









	
  

 	
  

 
	
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However,
for a moment by using this catalyzer it is difficult to achieve the diameter of
SiNWs less that 200 nm. The electronic states in the bandgap may also come from
a lack of passivation of surface defects. The passivation step is rather
crucial for SiNW solar cells, since SiNW have very high SVR ratio and their
opto-electronic properties strongly depends on the surface passivation.





Fig.
31. (a) Structure of the SiNW array solar cell. A p-n junction is formed
between the n-type SiNWs and the p-type Si substrate; (b) Dark and illuminated
I-V measurements of n-type SiNWs on p-type Si substrate (Perraud et al., 2009).


The
theoretical value of the efficiency for Si nanowire solar cells is predicted to
be as high as 16%, which makes them perfect candidates for higher bandgap
bricks in all-Si tandem cell approach. The first prototypes of SiNW solar cells
have excellent antireflection capabilities and shown the presence of the
photovoltaic effect. However, up today there was no evidence that this photovoltaic
effect occurred in a material with an increased bandgap.


5. Conclusions


Silicon
based third generation photovoltaics is a quickly developing field, which
integrates the knowledge from material science and photovoltaics. Today the
first prototypes of both Si QD solar cells and Si NWs solar cells have already
been developed. For a moment they present VOC, ISC and FF
values which still lower than those ones of the 1st generation PV
cells based on bulk Si – but all these problems are being addressed. It is too
prematurely to draw the conclusions while the further optimization steps of the
fabrication parameters were not performed. We should not forget that, for
example, although the airplane was not invented until the early 20th
century, Leonardo da Vinci sketched a flying machine four centuries earlier.


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