UNITED STATES


 


 


 


ARC PHOTOVOLTAICS CENTRE OF EXCELLENCE 2010/11 ANNUAL

REPORT, pp 61-99

ARC PHOTOVOLTAICSCENTRE OF EXCELLENCE ANNUAL REPORT,

2012, pp 169-224


 


 


 


Article referenced as support for the following
sections:


 


Page 8 & 9: Figure and Text on Tandem Solar Cells













4.5 Third Generation Strand – Advanced
Concepts


	
  

 
	
 University Staff:

 
	
 Prof.
 Gavin Conibeer (group leader)

 
	
 Dr Richard Corkish

 
	
 Prof. Martin Green

 
	
  

 
	
 Senior Lecturer:

 
	
 Dr Santosh Shrestha

 
	
  

 
	
 Senior Research Fellows:

 
	
 Dr Dirk KÖnig

 
	
  

 
	
 Lecturers:

 
	
 Dr Ivan Perez-Wurfl

 
	
  

 
	
 Research Fellows:

 
	
 Dr Shujuan Huang

 
	
 Dr Supriya Pillai (ASI
 Fellow)

 
	
 Dr Xiaojing “Jeana” Hao
 (ASI Fellow)

 
	
  

 
	
 Postdoctoral Fellows:

 
	
 Dr Pasquale Aliberti
 (until June 2012)

 
	
 Dr Robert Patterson (ASI
 Fellow)

 
	
 Dr Binesh Puthen-Veettil
 (ASI Fellow)

 
	
 Dr Xiaoming Wen (from Dec
 2012)

 
	
 Dawei Di (March to Sept
 2012)

 
	
 Craig
 Johnson (from Aug 2012) (ASI Fellow)

 
	
 Dr Siva Karaturi (from Feb
 2013)

 
	
  

 
	
 Professional officers:

 
	
 Dr Tom Puzzer (part time)

 
	
 Dr Patrick Campbell
 (shared with Thin Film)

 
	
 Mark Griffin (shared with
 Thin Film)

 
	
  

 
	
 Research Associates:

 
	
 Dr
 Didier Debuf (adjunct fellow)

 
	
 Yidan Huang

 
	
  

 
	
 Higher Degree Students:

 
	
 Andy Hsieh (until Nov
 2012)

 
	
 Dawei Di (until March
 2012)

 
	
 Craig
 Johnson (until Aug 2012)

 
	
 Sammy Lee (until Aug 2012)

 
	
 Haixiang Zhang

 
	
 Tian Zhang

 
	
 Yu “Jack” Feng

 
	
 Pengfei Zhang

 
	
 Sanghun Woo

 
	
 Yao Yao

 
	
 Neeti Gupta

 
	
 Chien-Jen
 “Terry” Yang

 
	
 Ibraheem Al Mansouri

 
	
 Yuanxun “Steven” Liao

 
	
 Suntrana Smyth

 


- 169 -









	
  

 
	
 Lingfeng Wu (from Mar
 2012)

 
	
 Xuguang Jia (from Mar
 2012)

 
	
 Ziyun Lin (from Mar 2012)

 
	
 Hongze Xia (from Mar 2012)

 
	
 Xi Dai (from Mar 2012)

 
	
 Shu Lin (from Aug 2012)

 
	
 Pei Wang (from Aug 2012)

 
	
 Simon Chung (from Aug
 2012)

 
	
  

 
	
 Casual Staff:

 
	
 James Rudd (part time)

 
	
 Adrian
 Shi (Aug to Dec 2012)

 
	
  

 
	
 Visiting Researchers:

 
	
 Dr Yukiko Kamikawa (AIST,
 Tsukuba, Japan, until Mar 2012)

 
	
  

 
	
 Undergraduate Students:

 
	
 4th year Thesis Projects:

 
	
 Dongchen Lan

 
	
 Guillaume Pauvert

 
	
 Aden Brown

 
	
 Li Shen

 
	
 Alexander Robinson

 



Abstract


There has been a significant increase in the work on hot
carrier cells in 2012 with the start of both the Australian Solar Institute (ASI) project and
US-Australia Solar Energy Collaboration (USASEC) projects. The Group IV nanostructure tandem cells project - the
“all-Si” tandem cell – has also seen an increase with the start of its own
ASI project. There has also been further work on up-conversion,
photo-electrolysis and on plasmonics.


There
has been increased understanding of the parameters required to grow good Si
nanostructure materials. Characterisation
of materials is leading to better understanding of the defects which narrow the
effective bandgap and hence limit VOC. Improvements in VOC
have been seen with lower temperature
measurements and in current with light trapping by various techniques.
Modelling using equivalent circuits is giving greater insight to device
performance.


Hot carrier cells efficiencies are being modelled with
greater input of real material properties and with integration of more complex structures.
Combined contact and absorber materials mdoels and those including both
multiple quantum well and phononic bad gap material have been developed. New
designs of contacts for rectification or for incorporation of energy selective
contacts directly in absorbers are being
developed. A wide range of potential materials for phononic modulated
structures have been identified and are being developed. Growth of
nanoparticle ordered arrays has progressed with
modelling of such superstructures also being further developed. Fully
integrated devices are now being designed and are expected to be
fabricated with collaborators in the near future.


The
up-conversion project now has good synthesis routes for Er containing phosphors
and for PbS downshifting quantum dots (QDs) used to sensitise Er to a wider
wavelength range. Progress on incorporating such materials with light trapping
structures is expected to give improvements in up-conversion for Si cells.


This
progress in all the main Third Generation project areas is improving
understanding and allowing optimisation of modelling, structures and devices.
Developments to come in the next year will see significant advancement in these
areas, with excellent prospects for good demonstration devices.


- 170 -






4.5.1 Third Generation Photovoltaics


The
“Third Generation” photovoltaic approach is to achieve high efficiency whilst
still using “thin film” second generation deposition methods. The concept is to
do this with only a small increase in areal costs and to use abundant and
non-toxic materials and hence reduce the cost per Watt peak [4.5.1]. Thus these
“third generation” technologies will be compatible with large scale
implementation of photovoltaics. The aim is to decrease costs to well below
US$0.50/W, towards US$0.20/W or better, by dramatically increasing efficiencies
but maintaining the economic and environmental
cost advantages of thin film deposition techniques (see Fig. 4.1.3 showing the
three PV generations) [4.5.1, 4.5.2]. To achieve such efficiency
improvements such devices aim to circumvent the Shockley-Queisser limit for
single band gap devices that limits efficiencies to the “Present limit”
indicated in Fig. 4.1.3 of either 31% or 41% (depending on concentration
ratio). This requires multiple energy threshold devices such as the tandem or
multi-colour solar cell. The Third Generation Strand is investigating several
approaches to achieve such multiple energy threshold device [4.5.1, 4.5.3].


The
two most important power loss mechanisms in single-band gap cells are the
inability to absorb photons with energy less than the band gap (1 in Fig.
4.5.1), and thermalisation of photon energy exceeding the band gap, (2 in Fig.
4.5.1). These two mechanisms alone amount to the loss of about half of the
incident solar energy in solar cell conversion to electricity. Multiple energy
threshold approaches can utilise some of this lost energy. Such approaches do
not in fact disprove the validity of the Shockley-Queisser limit, rather they
avoid it by the exploitation of more than one energy level for which the limit
does not apply. The limit which does apply is the thermodynamic limit shown in
Fig. 4.1.3, of 68.2% or 86.8% (again depending on concentration).





Figure 4.5.1: 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.


In
the Third Generation Strand, we aim to introduce multiple energy levels by
fabricating a tandem cell based on silicon and its oxides, nitrides and
carbides using reduced dimension silicon nanostructures to engineer the band
gap of an upper cell material. We are aiming to collect photo-generated
carriers before they thermalise in the “hot carrier” solar cell. Also we are
investigating absorption of two below bandgap photons to produce an electron-hole
pair in the cell by up-conversion in a layer behind the Si cell using erbium
doped host materials. In order to optimise the requisite properties, all these
structures are likely to be thin hence maximising absorption of light in thin
structures through light trapping is very important. Hence we are also
investigating localised surface plasmon enhanced coupling of light into these
Third Generation devices.


4.5.2 Si Nanostructure Solar Cells


Researchers: Shujuan Huang, Ivan Perez-Wurfl, Dirk KÖnig, Tom Puzzer, Xiaojing Hao,
Dawei Di, Zhenyu Wan, Sammy Lee, Yidan Huang, Lingfeng Wu, Xuguang Jia, Ziyun
Lin, Tian Zhang, Terry Yang, Gavin Conibeer, Martin Green


- 171 -






4.5.2.1 The ‘‘all-Si’’ tandem cell


We
are developing a material based on Si (or other group IV) quantum dot (QD) or
quantum well (QW) nanostructures, from which we can engineer a wider band gap
material to be used in tandem photovoltaic cell element(s) positioned above a
thin film bulk Si cell, see Fig. 4.5.2.


Previously
we have demonstrated the ability to fabricate materials which exhibit a blue
shift in the effective band gap as the QD or QW size is reduced, using
photouminescence [4.5.4] and absorption [4.5.5] data [4.5.6]. A thin film
deposition of a self-organised QD nanostructure is achieved through a sputtered
multi-layer of alternating Si rich material and stoichiometric dielectric
[4.5.4]. On annealing the excess Si precipitates into small nanocrystals which
are limited in size by the layer thickness, thus giving reasonable size uniformity,
as first demonstrated by Zacharias [4.5.7]. Demonstration of doping of these
layers with both phosphorus and boron to create a rectifying p-n junction has
resulted in devices with a photovoltaic open circuit voltage of 490mV [4.5.6,
4.5.7, 4.5.8].





Figure 4.5.2: A tandem photovoltaic cell using quantum confined QDs or QWs to
engineer the band gap of the top cell and potentially also the lower cells.
Short wavelength light is absorbed in the top cell and longer wavelengths in
lower cells, thus boosting the overall
voltage generated and hence efficiency. Formation of Si (or Ge or Sn)
QDs through layered thin film deposition of Si rich material which crystallises
into uniform sized QDs on annealing.


A
cell based entirely of Si, or other group IV elements, and their dielectric
compounds with other abundant elements (i.e. silicon oxide, nitride or carbide)
fabricated with thin film techniques, is advantageous
in terms of potential for large scale manufacturability and in long term
availability of its constituents. Such thin film implementation implies
low temperature deposition without melt processing, it hence also involves
imperfect crystallisation with high defect densities. Hence devices must be thin to limit recombination due to their
short diffusion lengths, which in turn means they must have high
absorption coefficients.


For
photovoltaic applications, nanocrystal materials may allow the fabrication of
higher band gap solar cells that can be used as tandem cell elements on top of
normal Si cells [4.5.11, 4.5.12]. For an AM1.5 solar spectrum the optimal band
gap of the top cell required to maximize conversion efficiency is 1.7 to 1.8eV
for a 2-cell tandem with a Si bottom cell [4.5.13]. To date, considerable work
has been done on the growth and characterization of Si nanocrystals embedded in
oxide [4.5.7, 4.5.14] and nitride [4.5.15, 4.5.16] dielectric matrices.
However, little has been reported on the experimental properties of Si
nanocrystals embedded in SiC matrix [4,5,17]. These are of particular interest
for application in photovoltaic devices because of an expected significant
increase in carrier transport due to a decrease in the barrier height between
adjacent nanocrystals [4.5.18]. As a result, sufficient carrier mobility can be
obtained to satisfy device fabrication requirements.


- 172 -






4.5.2.2 Fabrication of Si QD nanostructures


Thin
film techniques are used for nanostructure fabrication. These include
sputtering and plasma enhanced chemical vapour deposition (PECVD). The
deposition is a variation of the multi-layer alternating ‘stoichiometric
dielectric / Si rich dielectric’ process, shown in Fig. 4.5.3, followed by an
anneal during which Si nanocrystals precipitate limited in size by the Si rich
layer thickness [4.5.12, 4.5.7]. The most successful and hence most commonly
used technique is sputtering, because of its large amount of control over
deposition material, deposition rate and abruptness of layers. A multi-target
remote plasma sputtering machine with two independent RF power supplies as well
as additional DC power supplies is used in this work.





Figure 4.5.3: Multilayer deposition of alternating Si rich dielectric and
stoichimetric dielectric in layers of a few nm. On annealing the Si
precipitates out to form small nanocrystals of a size determined by the layer
thickness. Nanocrystal or quantum dot size is therefore uniform.


RF magnetron sputtering is used to deposit alternating
layers of SiO2 and SRO of thicknesses down to 2nm. [SRO refers to Si
rich oxide, formed by co-sputtering Si and SiO2.] Deposition of
multi-layers, consisting
typically of 20 to 50 bi-layers, is followed by an anneal in N2 from
1050 to 11500C. During the anneal, the excess silicon in the SRO layer
precipitates to form Si nanocrystals between the stoichiometric oxide layers.


For Si QDs in SiO2 the precipitation occurs
according to the following:


Precipitation
of excess Si from Si rich dielectrics in SiNx and SiC follows a
similar crystallisation reaction as Si precipitates from the amorphous matrix.
The techniques have also been applied to growth of Sn and Ge QDs in SiNx
in SiO2. Ge quantum dots can be precipitated at substantially lower
temperature, as discussed below.


4.5.2.3 Silicon nanocrystal devices on quartz substrates


Researchers: Ivan Perez-Wurfl, Xiaojing Hao, Dawei Dai, Adrian Shi, Lingfeng Wu,
Xuguang Jia, Ziyun Lin, Tian Zhang, Terry Yang


As
previously reported in previous annual reports, devices have been fabricated
using the SiQD in SiO2
materials, with p-n junction formation using B or P of multilayers,
respectively [4.5.9, 4.5.10, 4.5.19]. The fabricated p-n diodes
consisted of sputtered alternating layers of SiO2 and SRO onto
quartz substrates with in-situ boron and phosphorus 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 fabricated interdigitated solar cells have an effective area of up to
0.12cm2. The devices exhibit rectification and a photovoltaic
response with VOC up to 493 mV, but with as yet very small currents
and very bad fill factors. These are partly due to the very high


- 173 -






resistance
of the material and in particular to the relatively long lateral paths to
contacts at the back contact, necessitated by growth on an insulating quartz
substrate. The device structure with appropriate contacting is shown
schematically in Fig. 4.5.4. This also shows the current crowding effect which
results from the high lateral resistance.





Figure 4.5.4: Schematic diagram of the of the fabricated inter-digitated Si QD in
SiO2 devices, with B and P doping to create the p-n junctions and
illustrating the effect of current crowding due to high lateral resistance to
the back contact.


4.5.2.4 Equivalent circuit model of nanocrystal devices on
quartz substrates


In
order to better understand the limitation imposed by the device on the solar
cell performance, it is necessary to find a good equivalent circuit model
specific to our devices. The high resistivity of the base layer makes it
imperative to consider the two dimensional effects of current flow [4.5.20]. In
order to analyse the I-V characteristics
of these diodes we first generalised the model to include any number of diodes
in series. The series connection of diodes is used to explain the ideality
factors higher than two normally observed in our structures. We believe this is
a reasonable model as the ideality factor observed is almost independent of the
diode current. This type of behaviour has also been observed in multi quantum
well laser diodes [4.5.21], where it has been proven to arise from an
unintended series combination of diodes. Based on this series combination of
diodes, it is possible to obtain an expression relating temperature dependent I-V measurements to the band gap [4.5.9].
It is further possible to linearise the expression around an average measuring
temperature, Tavg, as
follows:





 (4.5.1)


where
k is Boltzmann’s constant, q is the electron charge, Tnom is the temperature at
which the saturation current, Ioi,
is defined, ni is the
diode ideality factor, and αi
is the saturation current temperature exponent. Notice that this equation
shows that the I-V characteristics
are related to a sum of band gaps.


As
the current flows from the base contact to the emitter, a linear voltage drop
along the base and under the diode isolation mesa causes an exponential change
in the diode current. This crowding of the current at the edge of the diode
mesa, depicted in Fig. 4.5.4, can be modelled adding a current dependence on
the series resistance. This series resistance, RS,
can be expressed as the sum of a current
independent, Rext and
a current dependent series resistance Rint
arising from current crowding:





 (4.5.2)


The
value of this resistance can be found by numerically solving the following
transcendental equation:


- 174 -









 (4.5.3)


Within
this mathematical framework it is then possible to extract the value of the
series resistance, remove its effect from the measured I-V characteristics to extract the actual
ideality of the diode current. We normally observe an ideality factor of 3.
Based on the proposed model of at least two diodes
in series, without any assumptions, it is only possible to establish that the
value of the band gap extracted from temperature dependent I-V measurements corresponds to a sum of
the band gaps (or activation energies) of these diodes.


An
interesting behaviour observed in these devices was the apparent lack of
correspondence between the dark and light I-V
characteristics. The series resistance extracted from the diode dark
I-V characteristic is too small
to explain the limited short circuit current as well as the low fill factor
measured under a simulated 1-sun condition. A more complete circuit model is
necessary to explain this discrepancy. We have proposed a model where the
observed behaviour is due to two distinct areas in the fabricated devices. The
photocurrent is produced only in a small area of the device, this area being
proportional to a fraction of the normalized diode area. In a Spice circuit
model this area is given a value smaller than 1, that we denote as fraction.
This will be a fitting parameter to reproduce the measured dark and illuminated
I-V characteristics. Figure 4.5.5
shows the circuit proposed.





Figure 4.5.5: Equivalent Spice circuit representation of the fabricated devices.


A
relatively large percentage of the device area is responsible for the measured
characteristics in the dark. The current in this area is caused by the
diffusion of minority carriers caused by the applied voltage (V1 in Fig. 4.5.5)
from the p or n side to the opposite region. This current is expected to be
large due to the low lifetime of the minority carriers (mostly recombination
current in the depletion region). The observed series resistance will be
proportional to the total series resistance, Rtot, and inversely proportional to the area, R2=Rtot/(1-fraction), where the diffusion current
occurs. The dark I-V behaviour
is modelled by the top branch of the circuit depicted in Fig. 4.5.5. D_QD1 and
D_Sch represent the series connection of diodes whose band gaps are extracted
using Eq. (4.5.1). The series resistance, R2, has the temperature and current
dependence detailed in Eqs. (4.5.2) and (4.5.3).


Only
a small part of the diode area may have a large enough lifetime to produce a
photocurrent. As this photocurrent flows only through this fractional area, the
series resistance is inversely proportional to
this fraction: R1=Rtot/fraction.
Since the photocurrent, Iph, flows through R1, the illuminated I-V characteristics are limited by
a larger resistance than that observed in the dark condition, as long as the fraction
of the diode area is smaller than one half. The simulations depicted in Fig.
4.5.6 show the reduction of ISC as the fraction, f, is varied from 99% to 1% of the total diode area.


- 175 -









Figure 4.5.6: Spice simulations of 1-sun I-V characteristics based on the circuit
depicted in Fig. 4.5.5. The fraction, f,
represents the normalized area of the diode where the photocurrent is produced.


With
the circuit model described, it will be possible to extract complementary
information from dark and illuminated I-V measurements.


In view of these simulations, it is clear that the
electrical characterisation of our devices needs to take into consideration previously overlooked
limitations. For example, great care should be taken when interpreting the
Quantum Efficiency extracted from a spectral response measurement as the
assumed condition of short circuit current may be incorrect even if the device
is externally short circuited (internally, the diode may be forward biased).
Moreover, as the current is proportional to the photon flux, and the flux is
generally different at each wavelength tested, the internal bias of the device
can be different at each point of the spectrum investigated.


4.5.2.5 Modelling for device optimisation


Researchers: Ivan Perez-Wurfl, Ziyun Lin, Tian Zhang


When
designing and implementing the Si-QD solar cell model the most important
performance limiting factor identified was the in-plane resistivity of the
material. To take this into consideration a 2-D model needed to be implemented.
This model takes into consideration the lateral flow of current in the solar
cell. By implementing this 2-D effect we have been able to recreate the
current-voltage (I-V)
characteristics of diodes made with Si-QD on insulating substrates. The model
takes into consideration the temperature dependence of the resistivity and
saturation current. The actual behaviour of the resistivity and saturation
current are modelled in terms of activation energies obtained from temperature dependent measurements. The activation energy
of the resistivity is related to defects in the material which act as
electronic traps. The activation energy of the saturation current is related to an effective electronic bandgap
that, in some cases, can be modified by the doping itself. In crystalline
semiconductors such as Si, doping does not affect the electronic bandgap unless
extreme levels of dopants are used (> 1x1019cm-3).
However, in disordered or amorphous semiconductors, doping the material will
always affect the electronic bandagap. For example, doping hydrogenated
amorphous silicon (a:Si:H) increases the defects in the material which in turn
act as traps. The end result is a highly resistive material despite the very
high doping levels used. Furthermore, the high doping itself adds such a high density of defects that the electronic
bandgap of the material is clearly reduced. To minimize the effect, undoped
materials are commonly used for the active PV layer of the device. We
have identified that the type of Si-QD films required for PV devices, has
indeed a similar behaviour as that of a-Si:H when it comes to doping. What is
more, the native defects in the Si-QD layers we produce are high enough to
affect the electronic bandgap even when no intentional doping is present.


- 176 -






The realisation that the effective electronic bandgap is
in fact the most important aspect limiting VOC has clarified an optimisation plan to follow.


Efforts
are now in place to establish optical methods to quantify the effective
electronic bandgap. Specifically, we are working on tailoring standard
characterisation methods, Photo Thermal Deflection Spectroscopy (PDS) and Photo
Luminescence (PL), to specifically study Si-QD materials. Both techniques offer
the advantage of allowing evaluation of defects. These defects should be
minimised to reap the benefits of the expected enhanced bandgap and, with the
aid of these techniques, could be
quantified without the need to fabricate devices. Once the quality of material
has been optimised according to minimised defects determined using these
optical methods, this material can then be used to fabricate devices.


4.5.2.6 Light trapping for increased current in devices


Researchers: Ivan Perez-Wurfl, Supriya Pillai, Lingfeng Wu, Ziyun Lin, Xuguang Jia


The photogenerated current in p-n unction Si
nanostructure devices has been increased by more than 100% compared to a cell
without any light trapping. The application of an Aluminium back-reflector gave a current enhancement of up to 46%.
Plasmonic silver nanoparticles gave an even more impressive result increasing
the current by up to 109%.


The
work concentrated around devices fabricated on quartz substrates in the
substrate configuration. This means that the devices were designed such that
illumination was applied from the top of the structure. The substrate is used
as a support for the device as well as the back-reflector as shown in Fig.
4.5.7.





Figure 4.5.7: Schematic of the structure used to optimise photocurrent using an Al
back-reflector.


The
amount of light absorbed by these structures is limited by the thinness of the
structure and the relatively low silicon content. Furthermore, light trapping
is not possible using a textured substrate as the material is deposited by
sputtering which is not as conformal as CVD or ALD, for example. Furthermore,
the thinness of the layer does not allow the typical chemical texturing used in
standard Si solar cells. The most straightforward way to increase light
absorption, and therefore short circuit current, is with the use of a metal
back-reflector. The efficiency of a back-reflector can be further enhanced (or
substituted) with the use of plasmonic effects that could increase the optical
absorption path by scattering incoming light.


A
sample was chosen such that the current was as high as possible for multiple
devices contained within the same insulating substrate. We will refer to this
sample as SM2A from here on. The devices showed a reasonable uniformity within
the sample to allow a statistically significant comparison between cells with
and without a back-reflector. The resistivity of the base had to be chosen to
be as low as possible to reduce its effect on the current without compromising
the quality of the diodes fabricated.


- 177 -






The
results of the photocurrent enhancement are summarised in Table 4.5.1. Notice
that the lowest enhancement is 31.3%. The average enhancement across the sample
was 40% with a maximum enhancement of 46.1%.


Table 4.5.1: Summary of
photocurrent enhancement achieved with an Aluminium back-reflector on sample SM2A.


	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 
	
  

 	
  

 	
 5mm2 device

 	
  

 	
 1mm2
 devices

 	
  

 
	
  

 	
  

 	 

 	
  

 	 

 	
  

 
	
  

 	
  

 	
 Device A

 	
  

 	
 Device B

 	
  

 	
 Device C

 	
  

 	
 Device D

 	
  

 
	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 



 	 

 
	
 ISC
 without Reflector (mA)

 	
  

 	
 18.7

 	
  

 	
  

 	
 4.1

 	
  

 	
  

 	
 4.6

 	
  

 	
  

 	
 4.6

 	
  

 	
  

 
	
 ISC with
 Reflector (mA)

 	
  

 	
 27.3

 	
  

 	
  

 	
 5.4

 	
  

 	
  

 	
 6.7

 	
  

 	
  

 	
 6.4

 	
  

 	
  

 
	
 % increase using Al (M2A)

 	
  

 	
 46.1

 	
 %

 	
  

 	
 31.3

 	
 %

 	
  

 	
 45.1

 	
 %

 	
  

 	
 37.5

 	
 %

 	
  

 





Figure 4.5.8: Typical dark and illuminated I-V characteristics
of a 1mm2 diode tested in sample SM2A. Three curves are shown:
Illuminated with Reflector (green), Illuminated with No Reflector (grey) and in
the Dark (blue).


Larger devices were also tested with the best result
shown on a 5mm2 device.


The largest device showed the highest percent
enhancement in photocurrent when tested with a back-reflector or with silver
nanoparticles. This particular enhancement is probably due to variations in the
device performance across the sample rather than having an actual dependence on
the area. It is worth pointing
out that the largest device is closer to the centre of the sample which
positively impacts the quality of the device. An interesting feature observed
in Fig. 4.5.8 is the larger VOC observed when tested without a back
reflector. This is probably due to heating of the sample while testing,
consequently, the temperature may not have been the same for all the
measurements. The actual temperature of the device is difficult to control due
to the thermally insulating substrate used. Details of the thermal management
issues will be discussed below.


Silver
nanoparticles have also been investigated. The enhancement obtained in the
photocurrent was significant. When tested
on a 5mm2 solar cell, the current improved more than two fold.
Specifically, the current was increased by 109% compared to that
obtained without any light trapping scheme.


The
silver nanoparticles were deposited on the top of the cell. Illumination was
then applied from the side of the substrate after removing the Al back
reflector as shown in the schematic in Fig. 4.5.9.


- 178 -









Figure 4.5.9: Schematic of the structure used to optimise photocurrent using silver
nanoparticles.


The
nanoparticles were obtained by evaporating 18nm of Silver followed by a 1 hour
anneal in a nitrogen purged oven at 200oC. Figure 4.5.10 shows the
comparison between the I-V characteristics
of the cell tested with an aluminium back-reflector or silver nanoparticles.





Figure 4.5.10: Illuminated I-V
characteristics of a 5mm2 diode showing an enhanced
photocurrent due to Silver nanoparticles (green curve) or an Al backreflector
(blue curves) compared to the same cell without any light trapping (red curve).


4.5.2.7 Devices on conductive substrate


We
have investigated three metals and one highly doped semiconductor thin film
suitable for use as conductive virtual substrates on a quartz supporting
substrate.


The
first all Si-QD devices demonstrated by our group were fabricated on quartz
substrates. The main reason for using quartz substrates is to demonstrate that
any PV response is in effect due to the Si-QD material and not the substrate
itself. The disadvantage of using these substrates is the current crowding that
occurs due to the high resistivity of the Si-QD layers. This current crowding
effect limits the maximum current that can be extracted from the solar cell.
Conductive substrates are investigated to mitigate this detrimental effect.


We
note that a thick substrate is not actually necessary. A thin, highly
conductive film is sufficient to reduce current crowding effects to a negligible
level. The conductive film of choice will be deposited on a quartz substrate to
ensure compatibility with our current process.


The high temperature needed to precipitate the Si-QD
from the silicon rich oxide layers, precludes the use of low melting point metals such as
Aluminium. To ensure compatibility with the unavoidable


- 179 -






high
temperature of our process, we chose three refractory metals with the highest
melting points, namely Tungsten (W),
Tantalum (Ta) and Molybdenum (Mo). We also investigated the use of silicon rich
silicon carbide (SRC) layers. SRC can be deposited with a low resistivity and
is perfectly compatible with the sputtering tool used for depositing SRO
materials.


4.5.2.7.1 Metals as back contacts


Researchers: Ziyun Lin, Ivan Perez-Wurfl, Xuguang Jia, Terry Yang


Sputtering
is the most common technique used to deposit refractory metals. The deposition
itself is straightforward but much optimisation is needed to ensure the stress
in the films is low enough to avoid the films delaminating from the substrate.
Early on in the optimisation process we identified severe difficulties in obtaining good quality films with W and Ta. We
concentrated our efforts then on improving the quality of Mo thin films.
It was found that a good indicator of the compatibility of the film with a high
temperature process was the minimisation of pinholes in the film. To accurately
quantify the density of pinholes we developed a technique based on the optical
transmission of Mo thin films. We have been able to use this characterisation
technique to quantify the pinhole density down a 1/100,000 total area of the
film. Fig. 4.5.11 shows the measured optical transmittance of two Mo films
deposited on quartz.





Figure 4.5.11: Optical transmission of two Mo films on quartz.


The
transmittance was measured using a Perkin Elmer Lambda 1050 spectrophotometer.
The measurement has been optimised by
ensuring a baseline calibration compatible with measurements of very low
transmittance. Figure 4.5.11 shows the nominally zero transmittance (10-6
%) of a 3mm thick Al slab. This “zero” level shows a clear differentiation with
a film transmitting only 0.001%. This optical characterisation method is fast
and accurate and allows us to pre-screen Mo films before using them as a substrate.


The
resistivity, surface roughness and the areal density of pinholes of the
optimised Mo layer is reported in Table 4.5.2.


Table 4.5.2: Characteristics of the best overall Mo film used as a substrate for a
Si-QD p-i-n structure.


	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 	
  

 
	
 Substrate 

 deposition 

 temperature (oC)

 	
  

 	
  

 	
 Film 

 thickness (nm)

 	
  

 	
  

 	
 Surface
 roughness (nm)

 	
  

 	
  

 	
 Sheet
 resistance 

 (W/square)

 	
  

 	
  

 	
 Area of pinholes 

 (% of total area)

 	
  

 
	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 	 

 
	
 420

 	
  

 	
  

 	
 350
 ± 50

 	
  

 	
  

 	
 0.84
 ± 0.2

 	
  

 	
  

 	
 2.35
 ± 0.8

 	
  

 	
  

 	
 0.0018

 	
  

 


- 180 -






A 300 nm Mo film similar to that of Table 4.5.2 but with
50X higher pinhole density was sputtered on quartz at room temperature. A bilayered
structure consisting of 25 bilayers (4nm SRO:4nmSiO2), nominally
un-doped with 45% Si by Volume in the SRO, was simultaneously deposited on a
quartz substrate and the Mo film. Both structures were also simultaneously
annealed at 1100oC for 1 hour in a nitrogen purged quartz tube
furnace. Visually, there was no apparent difference between the structure
deposited on Mo and that deposited on quartz after annealing. Comparing the PL
emission from the two samples also showed no difference as shown in Fig.
4.5.12.





Figure 4.5.12: a) Picture of the annealed 25 bilayer
structure on Mo b) Comparison of the PL emission from the multilayer on quartz
and Mo.


This
experiment proved that Mo is indeed compatible with the high temperature
process and has no obvious interaction with the multilayer structure.


The
same Mo layer was then used as a back contact for a standard full p-i-n
structure (70 bilayers total, bilayer = 4nm SRO/2nm SiO2, SRO = 66%
Si by volume). After the annealing process it was clear that Mo had interacted
with the p-i-n structure as can be seen in Fig. 4.5.13.





Figure 4.5.13: Picture of p-i-n structure after 11000C
anneal. The bottom left corner shows clearly that a reaction occured between
the p-i-n structure and the underlying Mo thin film.


Electrical
measurements showed that a dead short was created across the full p-i-n
structure precluding the measurement of VOC
or ISC. It is very likely that the formation of a Silicide caused
the shunt observed across the diode. The silicide formation is very
sensitive to the presence of SiO2 between Si and Mo. The thicker
oxide and lower Si content of the 25 bilayers may have been enough to avoid a
reaction with the underlying Mo thin film. In the case of the p-i-n structure,
the high concentration of Si in the SRO and the thin SiO2 barriers
promote the formation of a Silicide. Mo films
are still under investigation. Further optimisation of these films is under
intensive study. The use of a silicide is being considered to limit the
reaction of the metal and the p-i-n structure. Our findings in this critical
area will be reported in future reports.


- 181 -






4.5.2.7.2 Wide bandgap semiconductor as back contacts


Researchers: Dawei Di, Ivan Perez-Wurfl, Gavin Conibeer


A
second option investigated, geared at reducing the lateral resistivity, is the
use of a thin, highly doped, wide bangap material.


Silicon
rich carbide (SRC) layers have been investigated as the back contact for a
p-i-n structure. A significant advantage of using SRC rather than a metal
conductive substrate is the large amount of transmission
of SRC as compared to zero transmission for metals. This allows a wider range
of device testing including rear illumination and is more compatible
with a full tandem cell structure.


The
SRC layers are deposited directly on quartz. The films are doped with
phosphorous. After 1100oC annealing for 1 hour, the SRC layers show
a sheet resistance an order of magnitude smaller than that of bilayered SRO
layers of similar thickness and with the highest silicon content of interest
(66% Si by Volume).


Despite
its greater transmission than metals SRC as a conductive substrate nonetheless
does have an unavoidable absorption loss caused by the layer itself. The SRC
layer has to be as thin as possible to reduce
this loss. In the interest of investigating the effectiveness of SRC as a
conductive substrate, we choose a relatively thick layer that would be
easier to handle and characterise. With enough evidence to prove the
effectiveness of SRC we can then optimise the thickness without significantly
sacrificing its conductivity.


A
237 nm thick phosphorous doped SRC layer was deposited on quartz, annealed and characterised
prior to using it as a substrate. The measured sheet resistance of this film
was 5 k /square. The layer was then used as a substrate for a standard p-i-n
structure. The full stack was annealed once more, devices were fabricated and
characterised. It was possible to demonstrate a working device on this structure. Dark and light I-V characteristics of a 1mm2
diode is shown in Fig. 4.5.14.





Figure 4.5.14: I-V characteristics
of a standard p-i-n structure on and SRC substrate.


The
measured ISC and VOC are lower than those of a standard
p-i-n structure deposited directly on quartz. The reason for this lower
performance is mainly due to the absorption of light in the SRC layer which is
most probably not converted into current.


- 182 -






4.5.2.8 Increase in open circuit voltage


Researchers: Lingfeng Wu, Xuguang Jia, Ziyun Lin, Ivan Perez-Wurfl


The
original intention to show a path towards improving the open circuit voltage of
these devices beyond 493mV (the previous record voltage) was to vary the silicon
content in the SRO in order to exploit a suspected trend relating it to the
open circuit voltage. Our previous experience had shown that by reducing the
silicon content it is possible to increase the open circuit voltage sacrificing
some of the short circuit current due to
the inherent increase in the resistivity of the material. This suspected trend
had to be revised after developing the proper techniques to reproduce material
with the same silicon content on a regular basis. Contrary to expectations, the
change in silicon content of the SRO has very little effect on the open circuit
voltage.


An open circuit voltage of 503mV was measured on a 1mm2
solar cell at 250 K under an approximate 0.8 suns illumination condition. The demonstration of a voltage higher
than 500mV was one of the outcomes of this more structured line of work. The
reason for these non-standard testing conditions will be explained in detail in
the following paragraphs. It is important to point out that the demonstration
of a voltage larger than 500mV was measured on a device fabricated on an
oxidised Si wafer. Such devices can absorb less than half the number of photons
compared to a similar cell on a transparent quartz substrate. The I-V curve of this solar cell is shown in
Fig. 4.5.15.





Figure 4.5.15: Dark and illuminated I-V of cell with VOC = 503 mV,
under an approximate 0.8 suns, 250
K.


Having
developed the desired control on the composition and thickness, we deposited
p-i-n structures with variable silicon content (65.8, 63.7 and 60.9% Si by
Vol.) with nominally the same size of Si-QD. Using photoluminescence (PL) as a
metric to compare the similarity of QD in three different structures, we
confirmed that indeed the Si-QD were very similar in the three cases. The PL
spectra measured on these structures is shown in Fig. 4.5.16.


- 183 -









Figure 4.5.16: Measured PL on three p-i-n structures with
variable silicon content in the SRO (Si percent by Volume: 65.8%, yellow;
63.7%, pink; 60.9% blue).


As
can be seen in Fig. 4.5.16, the peak emission of the three structures is nearly
equal. The obvious differences are a small red shift of the PL emission spectra
for the 65.8% Si by Vol. sample and a slightly higher tail in the high energy
emission for the sample with the lowest Si content. The low variation in the silicon content was intentional
as these 5% variation represents a change in the lateral resistivity of
four to five orders of magnitude. The detailed relationship between Si content
and resistivity is currently under study.


The measured optical absorption of these three samples
confirms that indeed, the sample with a higher silicon content has the highest optical
absorption becoming slightly less pronounced as the silicon content is
decreased.


The fabricated devices expected to highlight the
relationship between VOC and Si content in the SRO, do not in fact show any particular trend. The
open circuit voltage measured ranges between 200mV and 300mV without any
apparent correlation to the silicon content. The lowest silicon content investigated
in fact did not produce any measurable devices, probably due to the extreme
high series resistance which in turn limits the device performance.


Having determined that the silicon content is not indeed
determining the open circuit of these devices, we set ourselves to the task of determining
the discrepancy of these observations with our previous experience showing an
actual dependence. We have concluded that the lack of accurate control in the
material composition in the early days of this research at UNSW, along with the
variation in the deposition rate, lead to unintended variations in the
deposited material which was interpreted as a trend
in VOC as a function of Si content in the SRO. We believe that there
may much more important factors affecting VOC besides the Si
content. Investigation of other options to increase VOC led us to
consider devices not fabricated on our standard quartz substrates.


Our
strongest proof that the substrates indeed affect the quality of the devices
built on them, come from an experiment performed by depositing single SRO
layers on quartz substrates. An apparent random
variation in obtaining a PL signal from only a handful of the samples prepared
was eventually correlated to an absorption peak observed in the IR
region of the spectrum, present only in the substrates that showed no PL
signal. This absorption peak (2.7mm) was later identified as an indicator of
OH groups. The higher concentration of OH groups caused out-diffusion of oxygen
from the substrate which oxidised the single layer deposited on them. The
vendor of the quartz substrates, however, does not differentiate between the
substrates with and without the OH related peak. In the same manner, there are
other impurities that could be infiltrating the Si-QD structures which could
potentially limit their electronic and PV quality.


- 184 -







With
this in mind we tested a p-i-n structure with an intended 64% Si by Volume in
the SRO layers. The structure was deposited on an oxidised Si wafer instead of
quartz. The oxide on the wafer is 97nm thick and ensures that there is no
possible electrical contact between the substrate and the Si-QD devices.


The
open circuit voltage measured in seven devices under 1sun illumination varied
between 420 mV and 440 mV. However, these devices are thermally isolated due to
the thick oxide present on the substrate. As the device is illuminated, it
heats up due to the light absorbed in the film however, due to the thermal
insulation caused by the substrate, a temperature gradient is bound to exist
with the highest temperature always being on the side of the device. An
advantage of measuring the I-V properties of these devices as a function of
temperature, is the possibility of determining the electronic bandgap of
the material used in the devices. The uncertainty in the actual temperature of
the device results in an uncertainty in the
determination of this bandgap. Fig. 4.5.17 presents a plot of VOC
vs. temperature measured on a 1mm2 Si-QD solar cell on a thermally
oxidised Si wafer.





Figure 4.5.17: VOC vs. Temperature measured on a
1mm2 Si-QD solar cell on a thermally oxidised Si wafer. The bandgap
extracted depends on the actual temperature of the device measured.


The
bandgap can be estimated based on the intercept on the y-axis of the linear fit
to the measured data. This intercept is
then the bandgap plus a correction factor in the order of 3kTaverage.
Based on this approach, we conclude that the effective electronic bangap
of the Si-QD layer is somewhere between 0.84 and 0.91 eV. These values are well
below the bandgap estimated from the PL peak emission reported in Fig. 4.5.16
(1.4 eV). The large discrepancy can be explained if the Si-QD material has a
very high density of defects bellow and above the conduction and valence band
edges respectively. We note that a lower electronic bandgap than anticipated
will make it difficult to obtain open circuit voltages as large as one would
expect from a material with a bandgap given by the peak of the PL emission.
Having identified this issue we can now concentrate on reducing the density of
electronic defects rather than tackling the resistivity or varying the diameter
of the Si-QD.


4.5.2.9 Other Si nanostructure materials studies


The use of matrices other than SiO2 can give
advantages in carrier transport. Or the use of NCs other than Si can give lower processing
temperatures and different possibilities for bandgap engineering.


- 185 -






4.5.2.9.1 Carrier tunnelling transport in Si QD superlattices


Transport
properties are expected to depend on the matrix in which the silicon quantum
dots are embedded. As shown in Fig. 4.5.18 different matrices produce different
transport barriers between the Si dot and the matrix, with tunnelling
probability heavily dependent on the height of this barrier. Si3N4
and SiC give lower barriers than SiO2 allowing larger dot spacing
for a given tunnelling current.





Figure 4.5.18: Bulk band alignments between silicon and its carbide, nitride and
oxide. Tunnelling probability between QDs separated by d depends exponentially on the square root
of the barrier height (ΔE1/2) multiplied by d. (e.g. [4.5.22] p244)


The results suggest that dots in a 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. Fluctuations in spacing and size of the dots can be
investigated using similar calculations. It is also found that the calculated
Bloch mobilities do not depend strongly on variations in the dot spacing but do
depend strongly on dot size within the QD material [4.5.18]. Hence, transport
between dots can be significantly increased by using alternative matrices with
a lower barrier height, ∆E.
For the same tunnelling current the spacing of QDs can increase for oxide to
nitride to carbide matrix.


4.5.2.9.2 Hetero-Interlayers


Researchers: Dawei Di, Zhenyu Wan, Gavin Conibeer


Experiments
have been carried out in which the layer between the Si rich layers is a
different dielectric. Si3N4 (nominally stoichiometric)
has two advantages. The higher density of the nitride restricts diffusion of Si
to nucleating nanocrystals during annealing. Thus size control and a more
mono-disperse size distribution of the nanocrystals result. Secondly the lower
barrier height of the nitride, as compared to the oxide within the Si rich
layers, leads to a higher tunneling probability in the growth direction, whilst
good quantum confinement is maintained in the nanocrystals planes.


Initial
results on this approach show promising decreases in the vertical resistivity
of such Si QD nanostructures with SiNx interlayers [4.5.23]. Figure
4.5.19 shows I-V curves measured on vertical samples on a conducting Si wafer
substrate, for both SiNx or SiO2 interlayers, each doped
with either B or P. In each case the substrate used was of the same carrier
type as the nominal type of the doped layer
on top. The most striking feature of the data is the much lower resistivity of
the B doped nitride interlayer samples as compared to the B doped oxide
interlayer sample. Whilst for P doping the oxide and nitride interlayer samples
show almost no difference in resistivity.


Hence
there is some evidence for significantly reduced resistivity for nitride as
compared to oxide interlayers, for the B doped samples. But the similar
resistivities for the two interlayers for P doping is not as first expected and
requires further investigation. The explanation is related to the differing
effects on Si QD crystallisation of P and B.


- 186 -









Figure 4.5.19: Logarithmic I-V curves
for Si QD samples grown with either nitride or oxide interlayers, and
subsequently annealed in H2 forming gas at 3500C. (a) B doped
samples: the nitride interlayer shows a dramatically lower resistivity; and (b)
P doped samples: the resistivities of the nitride and oxide interlayer samples
are similar. Insets show schematic diagrams of QD sizes and spacings.


Figure
4.5.20 shows the XRD spectra for SiO2 and SiNx
interlayers between SRO layers. Figure 4.5.20(a)
shows that for the oxide interlayer (i.e. for the SiO2-SRO-SiO2
structure) the effect of P (B) increasing (decreasing) the Si QD size is
demonstrated in the narrowing (broadening) of the XRD peaks for all of the
{111}, {220} and {311} Si peaks. The fact that all these principal peaks are
narrowed (broadened) to the same extent for a given sample indicates that the
effect is due to an increase (decrease) in QD size allowing a narrower
(broader) range of angles that satisfy the Bragg condition. This is in contrast
to inhomogeneous strain or a disordered structure in the nanocrystals which, if
it were present, would distort lattice spacing differently for different plane
families. Hence these QD sizes can be estimated using the Scherer equation for
peak broadening to give 5.7nm and 14.8nm for B and P, respectively [4.5.23].
Thus indicating that P incorporation significantly increases the size of Si
QDs. But in Figure 4.5.20(b) for the nitride interlayer the QD sizes are much
closer at 3.3nm and 5.0nm for B and P, respectively. This suggests that the
nitride interlayers suppress the growth of larger (smaller) nanocrystals for P
(B) doping respectively.





Figure 4.5.20: XRD comparison of P and B doped multilayers
for both (a) SiO2 and (b) SiNx interlayers between SRO
layers, all on quartz substrate.


The
likely reason for this uniformity of QD sizes for the nitride interlayer with P
and B doping, as compared to the oxide interlayer, is that the presence of a
strong gradient in N concentration at the interface between the silicon rich
oxide and silicon nitride layers will tend to promote heterogeneous nucleation
whether or not B or P are present. This will tend to suppress the differing QD
sizes in P and B material caused by their differing nucleation behaviour, thus
reducing the difference in QD spacing and its differing effects on resistivity.


This
suggests a reason for the differing I-V behavior shown in Fig. 4.5.19 based on
the opposite effects on Si QD crystallisation of P and B. For the oxide
interlayer samples the QDs are expected to


- 187 -






be
smaller with B and larger with P incorporation. The large size of the QDs in P
doped material are such that they penetrate several layers and significantly
reduce the number of tunneling events required in the vertical current
transport direction. This will tend to reduce resistivity. But in the nitride
interlayer the much more similar QD sizes for P and B doping shown in Fig.
4.5.20, mean this effect will be significantly reduced. Hence for B doping in
Fig. 4.5.19 (a) the QD sizes are very similar
and the presence of the lower barrier height nitride interlayer dominates in
reducing resistivity. Whilst for the P doping of Fig. 4.5.19 (b) the
larger QDs in the oxide interlayer material reduce resistivity to approximately
the same extent as does the lower barrier height of the nitride in the nitride
interlayer material even though this latter has larger spacing between the
smaller QDs.


4.5.2.9.3 Alternative group IV materials for tandem cells


Researchers: Sammy Lee, Shujuan Huang, Zhenyu Wan, Gavin Conibeer


Ge
nanocrystals in SiO2, made by a similar phase separation from solid
solution process, are being investigated as an alternative material with a
lower precipitation temperature and wider potential range of effective band gap
than Si nanocrystals. The temperature range employed is typically around 6500C,
but with substrate heating during growth can be as low as 3500C [4.5.24]. These
materials demonstrate a weak p-type conductivity without intentional doping and
can be made more p-type by addition of Sb during growth, although the exact
mechanism of doping is still under investigation.


Other
analogues involving Sn nanocrystals in dielectric matrices have been
fabricated. These require even lower processing temperatures, around 250-3000C,
and can be annealed in situ. However, Sn nanocrystals
formed in SiO2 matrix tend to undergo significant oxidation
[4.5.25]. In a Si3N4 matrix there is no oxidation
but the tetragonal β form of Sn always forms, because the β-Sn
allotrope is more thermodynamically stable than the cubic α-Sn allotrope
[4.5.26]. But as β-Sn is also semi-metallic as compared to the
semiconducting α-Sn, this also means that no semiconducting optical or
electronic properties of these Sn nanocrystal materials have yet been measured.


Other
alternatives to group IV elements in dielectric matrices are alloys of group IV
elements. Several of these are attractive
for controlling the band gap in a group IV tandem cell as summarized in [4.5.27].
SiC alloys with excess Si have been investigated both as material with
precipitated Si and SiC nanocrystals and as amorphous unannealed alloys. Band
gap control over the range of about 1.7 to 2.2eV is possible for the Si
nanocrystals material by control of nanocrystal size [4.5.28]. Si:Ge alloys are
of course well known as a means of continually varying the band gap from 0.7 to
1.1eV. But whilst they are suitable in a
tandem cell for cells under a Si cell, they are not suitable for a tandem cell
element on top of a Si cell.


Alloying
of Sn with small percentages of Ge can stabilize the cubic α-Sn phase
[4.5.29]. The Ge rich Ge:Sn alloy can also be grown. These alloys have band
gaps from 0.2-0.7eV and are hence potentially useful for cells to go under a Si
cell in a tandem, but again are not suitable for cells with higher band gap
than Si. However, the Ge:Sn alloy can also be used to stabilize the α
phase in nanocrystals incorporated in a dielectric matrix. SiGeSn alloys are a
variation of Sn:Ge alloys. The addition of less than 1% of Si allows lattice
matching of the alloy with Ge substrate [4.5.30] with addition of more Si also
allowing a higher range of direct band gap of 0.8-1.4eV [4.5.31]. The advantage
of the ternary alloy/compound is that the extra degree of freedom allows
independent control of band gap and lattice parameter, which in turn allows for
better quality materials.


Ge:C
and Sn:C alloys and GeC and SnC compounds are interesting as tandem cell
materials [4.5.27]. Ge:C offers a potentially wide band gap range from 0.6 to
1.1eV and a fairly continuous range of composition. Several of the material
properties of GeC have been calculated [4.5.32] and GeC thin films can be
sputtered [4.5.33] with material approaching stoichiometric GeC for high
methane gas flows. In addition high sputter powers lead to a greater degree of
sp3 hybridization of carbon bonds thus leading to cubic diamond or
zinc-blende structures as opposed to the graphitic structures for sp2 hybridization
[4.5.34].


- 188 -






4.5.2.10 Transfer to Si QD growth using PECVD


Researchers: Tian Zhang, Ivan Perez-Wurfl

Collaboration with: Manuel
Mogano, University of Milan; Birger Berghoff, RWTH Aachen


Sputtering
is a high energy process in which many defects are incorporated in deposited
material, it is also relatively slow for
thin film deposition. A transfer of the technology to plasma assisted chemical vapour
deposition (PECVD) has the potential advantages of improving material quality
for Si QD nanostructures in SiO2 matrix and of reducing long term
material costs.


PECVD
has been used to date primarily for growth of Si QD nanostructures in SiNx
and in SiC [4.5.35, 4.5.36]. It has now been extended to a SiO2
matrix.





Figure 4.5.21: XRR measurements of (a) as deposited PECVD
multilayers of SRO/SiO2 on quartz substrate (b) after annealing at
1100oC.


Figure 4.5.21 shows the X-ray reflectivity (XRR)
measurements of multilayer SRO/SiO2 material as deposited by PECVD and after annealing at
1100oC. They clearly show a good quality multilayered structure as
deposited with reduced quality after annealing. This is very consistent with
the formation of Si nanocrystals breaking up the multi-layer arrangement.





Figure 4.5.22: Raman spectroscopy of
annealed multilayer PECVD sample (blue line) compared to c-Si wafer.


Figure
4.5.22 shows Raman data on the annealed multilayer nanostructure. The clear
asymmetric broadening of the peak at around
510cm-1 strongly indicates formation of Si nanocrystals.


- 189 -






4.5.2.11 Summary of Group IV nanostructures for tandem cell
elements


Modelling
of the parameters of Si nanostrucrue growth and measured performance has
indicated that the band gap of these materials is dominated by defects. This is
allowing a more systematic investigation of the material quality in order to
move towards higher quality material to give higher VOC devices.
Currents have been improved through light trapping using both rear reflection
and plasmonics. On measuring devise significant temperature increase is seen
because of the high resistivity, therefore
lower temperature measurements of Voc have seen an increase to over
500mV for devices that are in effect locally at room temperature.
Reduced resistivity has been seen in structures with nitride interlayers and
other group IV elements have been incorporated into Ge QW and Ge:C alloys.
Further improvement of device parameters is expected as greater understanding
of the underlying material properties is gained with the current robust
characterisation methodology.


4.5.3 Hot Carrier Cells


Researchers: Shujuan Huang, Santosh
Shreshtha, Dirk KÖnig, Yukiko Kamikawa, Robert Patterson, Pasquale Aliberti, Binesh Puthen Veettil,
Ivan Perez-Wurfl, Andy Hsieh, Yu Feng, James Rudd, Pengfei Zhang, Yao Yao,
Hongze Xia, Neeti Gupta, Yuanxun Liao, Suntrana Smyth, Xi Dai, Pei Wang, Simon
Chung, Martin Green, Gavin Conibeer


Hot carrier solar cells offer the possibility of very
high efficiencies (limiting efficiency above 65% for unconcentrated
illumination) but with a structure that could be conceptually simple compared
to other very high efficiency PV devices – such as multi-junction monolithic
tandem cells. For this reason, the approach lends itself to ‘thin film’ deposition techniques, with their
attendant low costs in materials and energy usage and facility to use abundant,
non-toxic elements.


An
ideal hot carrier cell would absorb a wide range of photon energies and extract
a large fraction of the energy to give very
high efficiencies by extracting ‘hot’ carriers before they thermalise to the
band edges. Hence an important property of a hot carrier cell is to slow
the rate of carrier cooling to allow hot carriers to be collected whilst they
are still at elevated energies (“hot”), and thus allowing higher voltages to be
achieved from the cell and hence higher efficiency. A hot carrier cell must
also only allow extraction of carriers from the device through contacts which
accept only a very narrow range of energies (energy selective contacts or
ESCs). This is necessary in order to prevent cold carriers in the contact from cooling the hot carriers, i.e.
the increase in entropy on carrier extraction is minimized [4.5.37]. The
limiting efficiency for the hot carrier cell is over 65% at 1 sun and 85% at
maximum concentration – very close to the limits for an infinite number of
energy levels [4.5.1, 4.5.38, 4.5.39]. Fig. 4.5.23 is a schematic band diagram
of a hot carrier cell illustrating these two requirements.





Figure 4.5.23: Band diagram of the hot carrier cell. The
device has four stringent requirements: a) To absorb a wide range of photon energies;
b) to slow the rate of photogenerated carrier cooling in the absorber; c) To
extract these ‘hot carriers’ over a narrow range of energies, such that excess
carrier energy is not lost to the cold contacts; d) to allow efficient
renormalisation of carrier energy via carrier-carrier scattering.


- 190 -






Modelling
of Hot Carrier efficiencies has progressed with implementation of real material
properties to give more realistic efficiencies for InN which include Auger
processes and more realistic contact structures.
Significant progress has been made on demonstrating resonance in double barrier
selective energy structures. Further work on triple barrier double Si QW
structures has been carried out for rectifying ESCs. This is complemented by
improvements in 2/3D modelling of transport in these ESC structures. For
absorbers, modelling of nanocrystals superlattice arrays has been applied to
real material systems. The growth of such systems in both III-V QD
superlattices with collaborators and with colloidal Langmuir-Blodgett
dispersion of Si nanocrsystals has produced structures which are now being
characterised for their modulation of phononic properties. Also measurement of
carrier cooling rates has been extended to other large phononic gap bulk
materials including InN, demonstrating the importance of material quality.
Design of structures for hot carrier cells which should be practical and
realisable has developed, with the device properties more carefully specified
and plans for fabricating such structures in real devices.


4.5.3.1 Modelling of hot carrier solar cell efficiencies
using optimised absorber and ESCs


Researchers: Yu Feng, Pasquale Aliberti, Hongse Xia, Gavin Conibeer


The
performance of a hot carrier solar cell device has been evaluated taking into
account the non-equilibrium distributions of electrons and holes as well as
different rates of Coulomb scattering between charge carriers. This includes
processes of electron-electron (e-e) scattering, hole-hole (h-h) scattering and
electron-hole (e-h) scattering. Their effects on the carrier renormalization
and the electron-hole energy transfer have been quantitatively examined under
the relaxation time approximation (RTA). Other processes relevant to an
operating device, including photo-generation, radiative recombination, carrier
extraction, carrier cooling (inelastic carrier scattering) and impact
ionisation/Auger recombination (AR-II), have also been incorporated with RTA.


The
1000-sun performance has been predicted for the cell configuration of an intrinsic
InN absorber (thickness: 50nm) with two symmetric ESCs, in accordance with
previous work. The maximum efficiencies have been computed with different
relaxation time parameters, following the A and B given in the following
equation.


 





where
τe eq, τheq, τehehs are the
relaxation time of the respective processes (i.e. e-e/h-h/e-h scattering), N is
the carrier concentration.


The
results are shown in Fig. 4.5.24. In the upper figure efficiency decreases with
A, suggesting insufficient rates of e-e and h-h elastic scattering could reduce
the efficiency. This is because the carriers
with energies within the transmission window of the ESCs will be depleted
unless the elastic scattering is quick enough to supply new carriers
with these energies. The curve at B = 108 is almost flat, for e-h
elastic scattering quick enough to prevent carrier depletion, noting that e-h
scattering allows both carriers to exchange their energies, providing
possibilities of supplying carriers with the energies of extraction. For the
same reason, with negligible e-e and h-h scattering (A > 10-16)
the efficiency tends toward a constant value instead of dropping to zero. The
lower figure in Fig. 4.5.24 shows that the efficiency decreases with the
relaxation time of e-h scattering. It is partly due to the carrier depletion
especially with insufficient e-e and h-h scattering, at A = 10-14.
With sufficient e-e and h-h scattering, for example the equilibrium limit in
which carriers are completely re-normalised, the same trend still exists. In
fact a long relaxation time of e-h scattering blocks the energy flow from
electrons to holes, allowing the holes to relax rapidly due to their fast
thermalisation events compared to electrons
hence leaving the electron population hot. This would limit the hole flux out
of the device since only holes with an elevated energy and within a
small energy window could be extracted. On the contrary, the efficiency goes up
with B if we replace the hole ESC with a normal contact, one that


- 191 -






is
transmissive for all holes, since less e-h scattering means less energy loss
from hot electrons. To optimize the cell performance, the selection of the hole
contact depends on the e-h scattering rate in the absorber. A normal
non-selective valence band contact is beneficial for slow scattering rates (B
> 1011) while the ESC
configuration adopted here has a greater advantage for rapid scattering rates
(B < 1011).





Figure 4.5.24: Maximum efficiency with different A and B.
The e-h isolation limit indicates no scattering between electrons and holes.
The equilibrium limit indicates ultrafast e-e scattering and h-h scattering.


The
energy-dependent distributions of electrons and holes are shown in Fig. 4.5.25,
corresponding to the Maximum Power Points
(MPPs) with different A and B values. From left to right the figures show the
trend of increasing carrier temperature with e-e and h-h scattering rates, as
with less carrier depletion a faster
extraction rate can occur reducing the retention time of carriers. From top to
bottom the figures demonstrate a larger deviation between electron distribution
and hole distribution with less e-h scattering. The carrier depletion
within the transmission window of ESCs is illustrated in the corresponding
insets. The resupply of depleted holes is more rapid than electrons because of
the proportionality between the scattering rate and the carrier effective mass.
With a small A, i.e. fast e-e/h-h scattering rates, the supply of holes is
sufficient for the carrier extraction, leading to a larger current density and
a higher overall efficiency. In this case the electron depletion could be even
more severe since the increase in current depletes electrons more quickly. This
is observed with B = 1010 and 1012. Apart from the region
of carrier depletion the distributions are close to Fermi functions, which may
be used for defining a “quasi-temperature” for each type of carrier.


- 192 -









Figure 4.5.25: Energy-dependent distributions of electrons
(black solid line) and holes (red dashed line) when operating at the Maximum
Power Points (MPPs) with different A and B: A = 10-14 (left), A = 10-19 (right); B = 1010
(upper), B = 1012 (middle), B = 1014 (lower). The
depleted region within the transmission window is enlarged in the
corresponding inset.


The
calculated current-voltage curves are shown in Fig. 4.5.26. With faster e-h
scattering (smaller B) the short- circuit
current density is increased while the open-circuit voltage is reduced. It is
difficult to extract holes at the elevated energy level required by ESCs
unless e-h scattering is quick enough to resupply highly energetic holes. On
the other hand the increase of voltage with slower e-h scattering comes from
the hotter electron distribution at the open-circuit condition, with less
energy lost to the thermalised holes. If a
HCSC is designed with the narrow high energy transmission window for holes, as
adopted in this work, the maximum efficiency is always higher for quicker e-h
scattering. The drawback of slow e-h scattering would be pronounced for insufficient
e-e and h-h scattering rates, since the significant carrier depletion would
inhibit the carrier extraction resulting in a colder carrier distribution.
Hence the quasi-temperatures of both electrons and holes are low at MPPs, if
e-h interaction is weak, while in the lower inset the contrast between them
increases. The carrier quasi-temperature insets also indicate a test of the
validity of commonly used assumptions: electrons and holes share the same
temperature only for fast e-h scattering (B < 1010) while holes
are completely thermalized when e-h scattering rates are slow (B > 1014).


- 193 -






For
real materials the choice of parameters A and B can be determined by computing
the realistic scattering rates between carriers. We have calculated the scattering
relaxation times for bulk cubic-InN, taking into account both the direct
Coulomb potential and the exchange potential. The results closely follow the relation shown in the equation
above for relaxation times, with A = 2.5X10-19 and B = 2.4X1011.


The
results show that the maximum efficiency varies significantly with the carrier
scattering properties due to energy-dependent carrier depletion and asymmetric
statistics of electrons and holes. With different hole contacts the dominant
loss mechanism switches: with an ESC at an elevated energy the hole extraction
is blocked until significant heat is transferred from hot electrons. Thus at
the hole side an ESC is more suitable for fast e-h scattering while a normal
contact is beneficial for slow e-h scattering rates. In addition, the RTA model
suggested here provides the possibility of analyzing any realistic HCSC device.
It also serves to optimize the transmission properties of both contacts. According to this work it heavily
depends on the carrier scattering properties in the absorber.


4.5.3.3 Energy Selective Contacts (ESCs)


Researchers: Binesh Puthen-Veettil,
Yu Feng, Andy Hsieh, Yuanxun Liao, Santosh Shrestha, Gavin Conibeer


One
practical implementation of the requirement for a narrow range of contact energies
is an energy selective contact (ESC) based on double barrier resonant
tunnelling. Tunnelling to the confined energy levels in a quantum dot layer
embedded between two dielectric barrier layers, can give a conductance sharply
peaked at the line up of the Fermi level on the ‘hot’ absorber side of the
contact with the QD confined energy level. Conductance both below this energy
and above it should be very significantly
lower. This is the basis of the current work on double barrier resonant
tunnelling ESCs.


4.5.3.3.1 Modelling optimised materials for Energy Selective
Contacts


Energy selective contacts (ESCs) for a HC solar cell can
be implemented using different materials and structures. The main property that needs to
be achieved is precise energy selectivity in carrier extraction. Different
models have been developed to identify possible materials for ESCs and to optimize their properties. Amongst the different
possible material configurations, group IV and group III-V double
barrier resonant tunnelling structures look to be the most promising.





Figure 4.5.26: Schematic diagram of a
double barrier structure. Silicon QDs are formed in a dielectric matrix that may not be the same as the
barrier material. The growth direction of the structure is along the z axis.


Modelling
of group IV materials has shown that the structure with the highest potential
to be used as ESC consists of double barrier structure (DBS) with quantum dots
(QDs) embedded in a dielectric matrix.


Modelling
of electrical properties for the structure shown in Fig. 4.5.26 has been
performed for Si QDs in Si based dielectric materials. The reasons for focusing
on Si in the first instance for group IV


- 194 -






materials
are abundance and well established growth techniques. Structures consisting of
Si QDs and the following dielectric
barriers have been modelled: SiO2, Si3N4 or
SiC. Results show that optimal energy selectivity properties are
obtained with higher lateral confinement within the middle layer of the
structure and higher vertical conductivity. Amongst the material taken into
account in the simulations, structures consisting of Si QDs in a SiO2
with SiC barriers have shown the best overall energy extraction properties.


ESCs
for a III-V based HC solar cell can be realized using a QW structure in a DBS,
probably requiring either MBE or MOVPE growth. Modelling results have shown
that, for a HC cell based on an InN absorber, the optimum configuration for
ESCs is constituted by an InN/InXGa1-XN DBS [4.5.40].
This structure allows modulation of the extraction energy and shape of the
transmission probability function
independently for the electrons and holes contacts. The value of extraction
energy can be modified by engineering the stoichiometry and the
thickness of the QW structure. In general the optimal extraction energy for
electrons and holes is different, due to the different values of effective
mass. Thus, the hole ESC QW should be physically thinner and have a higher
barrier than the electron ESC QW to achieve similar currents at reasonable
extraction energies, as illustrated in Fig. 4.5.27.





Figure 4.5.27: Schematic of a HC solar cell based on a InN
absorber and InN/InGaN ESCs.


The
transmission coefficient of a double-barrier semiconductor structure has been
numerically calculated using a tight-binding method. The phonon scattering
during the transmission has been incorporated with a Greens function method.
This constructs a framework for simulating a real resonant tunneling system
with ultra-thin layers. It will help in optimizing the material system for the
ESCs. Some sampling results for the DBS of AlAs/GaAs/AlAs are shown in Fig.
4.5.28.





Figure 4.5.28: The total transmission (y axis: 01) profile
as a function of the normally-penetrating electron energy (x axis: eV) through
the zinc-blende DBS AlAs(3 unit cells)/GaAs (3 unit cells)/AlAs (3 unit cells),
with different phonon scattering properties. Left: No phonon scattering
considered; Right: with phonon scattering and the initial phonon occupation
number = 0.


- 195 -






Here
we mainly aim to utilize the 1st energy peak. From the transmission
calculation of an AlAs/GaAs/AlAs double-barrier structure, this peak is visible
and reasonably significant if only normally incident electrons are considered.
(These constitute the majority of energy transmission in any case.) With the
consideration of phonon scattering, side-peaks appear which are not favorable
for energy-selective extraction. This
indicates a preference of choosing less polar materials. Besides, the addition
of arbitrary directions off normal may broaden the transmission peak with a
little loss on the selectivity. Such a Greens function model can treat
any type of atomic system in terms of the scattering problem. It is ready for
use as a tool to choose the optimal material system for the energy-selective
transmission.


4.5.3.3.2 Triple barrier resonant tunnelling structures for
carrier selection and rectification


Researchers: Gavin Conibeer, Andy Hsieh, Santsoh Shrestha


A
normal energy selective contact (ESC) using double barrier resonant tunnelling
is symmetric and does not rectify electrons from holes. An ESC which uses
triple barrier resonant tunnelling through two separately tuned quantum well or
quantum dot layers can give a much greater energy selectivity and also is a
filter to separate electrons and holes in opposite directions (i.e.
rectification). The mini-band line-ups of the confined levels need to be tuned
by the well thicknesses, effective mass of electrons and holes and barrier
heights such as illustrated schematically in Fig. 4.5.29. With appropriate choice of these for the left contact
there can be an alignment of the confined levels such as to give a
resonant channel in the conduction band for electrons but misalignment of
levels in the valence band such that there is no channel for holes. Similarly
for the right contact the valence band confined
levels are aligned to allow holes to pass but misaligned in the conduction band
so as to block electrons. For non-equal electron and hole effective
masses (as is the case for most materials) this condition will in general be
met [4.5.41, 4.5.42].





Figure 4.5.29: Triple barrier double well resonant
tunnelling energy selective contacts. Wells are of different width designed to
give allignment of 1st and 2nd confined levels in the conduction
band on the left and for 1st and
2nd confined levels in the valence band on the right. In general for
non-equal effective masses allignment of first and second confined
energy levels for electrons will not give the same condition for holes thus giving
electron/hole rectification and a preferred collection direction in the
external circuit.


The collection channels for electrons and holes would
ideally be aligned with the peak occupancies of energy of the hot electron and hot hole
populations respectively, which in turn will depend on the respective
temperatures of these populations. This condition will give the highest
efficiency, but as discussed in Section 4.5.3.1, for high electron-electron and
hole-hole carrier scattering rates, it is not very critical.


- 196 -






Ideally
there would be two of these triple barrier double QW/QD devices on either side
of the absorber, but as there is much less energy in the high effective mass
hole population a simpler double barrier or simple band contact on the hole
side is also possible, without much loss in collected energy, as investigated
in Section 4.5.3.1. Various materials combinations could be used for such a
design, but specifically we are looking at III-V triple barrier QWs based on
the InGaN/GaN system. Again this is addressed to some extent in Section
4.5.3.1.


The
exact thickness of QWs or size of QDs and barrier thicknesses and contact work
functions will have to be tuned for the
particular confined energy levels involved in each material. Use of more than two
QW or QD layers on either side is also possible. This will increase energy
selection by requiring even more careful alignment.


This
is the first application of this concept to energy selective contacts for hot
carrier, Hot Lattice or Thermoelectric cells, but the concept itself was first
suggested by [4.5.43] and has been used for energy filtering in quantum cascade
lasers for several years. It is also not the first application of the concept to solar cells as [4.5.44], at Imperial
College, have used the concept for a multi-QW solar cell design. But it
is the first application to Hot Carrier or Hot Lattice devices and its ability
to rectify is likely to be a very significant advantage.


4.5.3.3 Hot Carrier Absorbers: slowing of carrier cooling


Carrier
cooling in a semiconductor proceeds predominantly by carriers scattering their
energy with optical phonons. This builds up a non-equilibrium ‘hot’ population
of optical phonons which, if it remains
hot, will drive a reverse reaction to re-heat the carrier population, thus
slowing further carrier cooling. Therefore the critical factor is the
mechanism by which these optical phonons decay into acoustic phonons, or heat
in the lattice. The principal mechanism by which this can occur is the Klemens mechanism, in which the optical phonon
decays into two acoustic phonons of half its energy and of equal and
opposite momenta [4.5.45]. The build-up of emitted optical phonons is strongly
peaked at zone centre both for compound semiconductor due to the FrÖlich
interaction and for elemental semiconductors due to the deformation potential
interaction. The strong coupling of the FrÖlich interaction also means that
high energy optical phonons are constrained to near zone centre even if
parabolicty of the bands is no longer valid [4.5.46]. This zone centre optical
phonon population determines the dominant optical phonon decay mechanism is
this pure Klemens decay.


4.5.3.3.1 Modelling of electron-phonon interactions:


Researchers: Yu Feng, Hongze Xia, Santosh Shrestha, Robert Patterson, Gavin Conibeer


Another
important factor is the Frohlich interactions between hot electrons and
longitude optical phonon modes. The
importance of investigating the fundamental particle interactions occurring in
the device is demonstrated by the significant dependence between the
carrier energy relaxation time and the output efficiency. The underlying
mechanism of phonon bottleneck effect involves two parts, the first is the
block of Frohlich interaction between electrons and phonons; while the second
is the block of phonon decay from the
optical modes into acoustic modes. Also the build-up of non-equilibrium of
acoustic phonons may also contribute to reduced carrier energy relaxations. In
order to understand the mechanism of slow relaxation rates in different
types of materials the three interaction processes need to be simulated.


The
Frohlich interaction occurs in polar semiconductors and creates energy exchange
between electrons and longitudinal phonons. It affects the total energy and
state transition rates by an interaction
Hamiltonian. The Frohlich interaction Hamiltonian for superlattice structures
involves an overlap integral over one dimension (the direction
perpendicular to planes). The derivation of the Hamiltonian involves the
interaction between electrons and polarization fields induced by longitudinal
mode vibrations. By obtaining the divergence of the polarization field an
effective phonon envelope function is developed and incorporated in the overlap
integral.


- 197 -






4.5.3.3.2 Phonon decay mechanisms


Researchers: Gavin Conibeer, Santosh Shrestha, Robert Patterson, Yu Feng, Hongze
Xia, Suntrana Smyth


An
optical phonon will decay into multiple lower energy phonons. The processes
require the conservation of energy and momentum. However, additionally, the
principal decay path must be into two LA
phonons only, which are of equal energy and opposite momenta, via an
anharmonicity in the lattice, this was suggested by Klemens [4.5.45] and
has been shown to apply to a wide range of materials.


Wide Phononic Gaps in III-Vs and analogues


For
some compounds in which there is a large difference in masses of the
constituent elements, there exists a large gap in the phonon dispersion between
acoustic and optical phonon energies. If large enough this phonon band gap can
prevent Klemens’ decay of optical phonons, because no allowed states at half
the LO phonon energy exist. InN is an example of such a material with a very
large phonon gap. The prevention of the Klemens’ mechanism forces optical
phonon decay via the next most likely, Ridley mechanism, of emission of one TO
and one low energy LA phonon. Such a mechanism only has appreciable energy loss
(although still much less than Klemens’ decay) if there is a wide range of
optical phonon energies at zone centre. This is only the case for lower
symmetry structures such as hexagonal. For a high symmetry cubic structure, LO
and TO modes are close to degenerate at
zone centre and the Ridley mechanism is severely restricted or forbidden.
Unfortunately cubic InN is very difficult to fabricate precisely because
of the large difference in masses that give it its interesting phononic
dispersion.


Slowed
cooling has been observed in some III-V compounds in which there is a large
difference in atomic mass. This has been shown in slowed carrier cooling for
InN [4.5.47] and for slowed carrier cooling in InP compared to the small mass
ratio GaAs [4.5.48].


Analogues of InN with abundant elements, but also with
narrow Eg include II-IV-VI compounds such as ZnSnN2, IIIA
nitrides such as LaN and YN which should have large phonon gaps, and IVA
nitrides such as ZrN and
HfN which are readily available. Bi and Sb compounds should also have large
phonon gaps but are not abundant. Group IV compounds have large calculated gaps
and small Egs, as well as several other advantages [4.5.49]. Fig. 4.5.30 shows
calculations of phonon dispersions for group IV compounds which have large
phonon band gaps, sufficient to block Klemens’ decay.





Figure 4.5.30: Adiabatic bond charge calculations of phonon dispersions for group IV
compounds. Phonon gaps increase as the mass ratio increases with those for GeC
and SnC large enough to block Klemens’ decay.


Slowed carrier cooling in MQWs


Low
dimensional multiple quantum well (MQW) systems have also been shown to have
lower carrier cooling rates. Comparison of bulk and MQW materials has shown
significantly slower carrier cooling in the latter. Fig. 4.5.31 shows data for
bulk GaAs as compared to MQW GaAs/AlGaAs


- 198 -






materials as measured using time resolved transient
absorption by Rosenwaks [4.5.50], recalculated to show effective carrier temperature as a
function of carrier lifetime by Guillemoles [4.5.51]. It clearly shows that the
carriers stay hotter for significantly longer times in the MQW samples, particularly
at the higher injection levels by 11⁄2 orders of magnitude. This is due to an
enhanced ‘phonon bottleneck’ in the MQWs allowing the threshold intensity at
which a certain ratio of LO phonon re-absorption to emission is reached which
allows maintenance of a hot carrier population, to be reached at a much lower
illumination level. More recent work on strain balanced InGaAs/GaAsP MQWs by
Hirst [4.5.52] has also shown carrier temperatures significantly above ambient,
as measured by PL. Increase in In content to make the wells deeper and to
reduce the degree of confinement is seen to increase the effective carrier
temperatures.


The
mechanisms for the reduced carrier cooling rate in these MQW systems are not
yet clear. However there are three effects that are likely to contribute. The
first is that in bulk material photogenerated hot carriers are free to diffuse
deeper into the material and hence to reduce the hot carrier concentration at a
given depth. This will also decrease the density of LO phonons emitted by hot
carriers as they cool and make a phonon bottleneck more difficult to achieve at
a given illumination intensity. Whereas in a MQW there are physical barriers to
the diffusion of hot carriers generated in a well and hence a much greater
local concentration of carriers and therefore also of emitted optical phonons.
Thus the phonon bottleneck condition is achieved at lower intensity.


The
second effect is that for the materials systems which show this slowed cooling,
there is very little or no overlap between the optical phonon energies of the
well and barrier materials. For instance the optical phonon energy ranges for
the GaAs wells and AlGaAs barriers used in [4.5.50] at 210-285meV and
280-350meV, respectively, exhibit very little overlap in energy, with zero
overlap for the zone centre LO phonon energies of 285 and 350meV [4.5.53].
Consequently the predominantly zone centre LO phonons emitted by carriers
cooling in the wells will be reflected from the interfaces and will remain
confined in the wells, thus enhancing the phonon bottleneck at a given
illumination intensity.


Thirdly,
if there is a coherent spacing between the nano-wells (as there is for these
MQW or superlattice systems) a coherent Bragg reflection of phonon modes can be
established which blocks certain phonon
energies perpendicular to the wells, opening up one dimensional phononic band
gaps (analogous to photonic band gaps in modulated refractive index
structures. For specific ranges of nano-well and barrier thickness these
forbidden energies can be at just those energies required for phonon decay. This coherent Bragg reflection
should have an even stronger effect than the incoherent scattering of the
second mechanism above at preventing emission of phonons and phonon decay in
the direction perpendicular to the nano-wells.


It
is likely that all three of these effects will reduce carrier cooling rates.
None depend on electronic quantum confinement and hence should be exhibited in
wells that are not thin enough to be quantized but are still quite thin
(perhaps termed ‘nano-wells’). In fact it may well be that the effects are
enhanced in such nano-wells as compared to full QWs due to the former’s greater
density of states and in particular their greater ratio of density of
electronic to phonon states which will enhance the phonon bottleneck for
emitted phonons. The fact that the deeper and hence less confined wells in
[4.5.52] show higher carrier temperatures is tentative evidence to support the
hypothesis that nano-wells without quantum confinement are all that are
required. Whilst several other effects might well be present in these MQW
systems, further work on variation of nano-well and barrier width and
comparison between material systems, will distinguish which of these reduced
carrier diffusion, phonon confinement or phonon folding mechanisms might be
dominant.


- 199 -









Figure 4.5.31: Effective carrier temperature as a function
of carrier lifetime for bulk GaAs as compared to GaAs/AlGaAs MQWs: time
resolved transient absorption data for different injection levels, from
Rosenwaks [4.5.50], recalculated by Guillemoles [4.5.51].


4.5.3.3.3 Hot carrier cell absorber requisite properties


The
above discussion allows us to estimate the major properties required for a good
hot carrier absorber material. These are listed below in approximate order of
priority, although their relative importance may well change in light of future
research:


	
  

 	
  

 
	
 1.

 	
 Large
 phononic band gap (EO(min) - ELA) - to suppress Klemens’ decay, requires
 large mass difference between elements.

 
	
  

 	
  

 
	
 2.

 	
 Narrow
 optical phonon energy dispersion (ELO - EO(min)) – to minimise the loss by
 Ridley decay, requires a high symmetry.

 
	
  

 	
  

 
	
 3.

 	
 Small Eg < 1eV – to allow broad range of
 photon absorption.

 
	
  

 	
  

 
	
 4.

 	
 A
 small LO optical phonon energy (ELO) – to reduce the amount of energy lost
 per LO phonon emission.

 
	
  

 	
  

 
	
 5.

 	
 A
 small maximum acoustic phonon energy (ELA). This maximises (EO(min) - ELA)
 and is important if ELO is also small.

 
	
  

 	
  

 
	
 6.

 	
 Good
 renormalisation rates in the material, i.e. good e-e and h-h scattering. This
 condition is met in most semiconductors quite easily, with e-e scattering
 rates of less than 100fs. But it may be comprised in nanostructures.

 
	
  

 	
  

 
	
 7.

 	
 Good
 carrier transport in order to allow transport of hot carriers to the
 contacts.

 
	
  

 	
  

 
	
 8.

 	
 Ability
 to make good quality, ordered, low defect material.

 
	
  

 	
  

 
	
 9.

 	
 Earth
 abundant and readily processable materials.

 
	
  

 	
  

 
	
 10.

 	
 No,
 or low, toxicity of elements, compounds and processes.

 


4.5.3.3.4 Hot carrier absorber: choice of materials


InN
has most of these properties, except 4, 8 & 9, and is therefore a good
model material for a hot carrier cell
absorber. InN is investigated below in combination with its alloys with GaN.
However the abundance of In is very low, so it is difficult to see InN as a
long term material suitable for large scale implementation. Hence analogues of
InN, which retain its interesting property of a large phonon band gap,
are also investigated.


Analogues of InN


As
InN is a model material, but has the problems of abundance and bad material
quality, another approach is to use analogues of InN to attempt to emulate its
near ideal properties. These analogues can be II-IV-nitride compounds, large
mass anion III-Vs, group IV compounds/alloys or nanostructures.


- 200 -









Figure 4.5.32: Use of the periodic table to analyse possible
analogue compounds of InN based on atomic mass combination and
electro-negativity.


II-IV-Vs: ZnSnN; ZnPbN; HgSnN;
HgPbN


With
reference to Fig. 4.5.32, it can be seen that replacement of In on the III
sub-lattice with II-IV compounds is analogous and is now quite widely being
investigated in the Cu2ZnSnS4 analogue to CuInS2
[4.5.54].


ZnGeN
can be fabricated [4.5.55] and is most directly analogous with Si and GaAs.
However, its band gap is large at 1.9eV. It also has a small calculated
phononic band gap [4.5.56]. ZnSnN has a smaller electronic gap (1 eV) and
larger calculated phononic gap [4.5.56]. It is however difficult to fabricate, and also its phononic gap is not as
large as the acoustic phonon energy making it difficult to block Klemens
decay completely. HgSnN or HgPbN should both have smaller Eg and
larger phononic gaps. These materials have not yet been fabricated [4.5.57].


Large mass cation:


The
Bi and Sb compounds have large predicted phononic gaps and Bi is a relatively
abundant material, with only low toxicity [4.5.57]. BiB has the largest
phononic gap but AlBi, Bi2S5, Bi2O3 (bismuthine)
are also attractive. Similarly SbB has a large predicted phononic gap. That for
AlSb is the same size as the acoustic phonon energy and its band gap is 1.5eV,
making it marginal as an absorber material and similar to InP.


Group IIIA III-Vs


LaN
and YN both have large phononic gaps whilst that for ScN is too small. The
Lanthanides can also form III-Vs. ErN and other RE nitrides can be grown by
MBE. The phononic band gaps of the Er compounds are predicted to be large,
because of the heavy Er cation, but its
discrete energy levels make it not useful as an absorber, although the
combination of properties in a nanostructure could be advantageous.


Group IV alloy/compounds:


All
of the combinations Si/Sn, Ge/C or Sn/C look attractive with large gaps
predicted in 1D models. However being all group IVs they only form weak
compounds. Unfortunately SiC, whether 3C, 4H or 6H, has too narrow a phononic
gap. Nonetheless GeC does form a compound and is of significant interest
[4.5.58].


- 201 -






There
are also several other inherent advantages of group IV compounds/alloys all of
which are associated with the four valence electrons of the group IVs which
result in predominantly covalent bonding:


	
  

 	
  

 
	
 a)

 	
 The
 elements form completely covalently bonded crystals primarily in a diamond
 structure (tetragonal is also possible as in βSn). However for group IV
 compounds the decreasing electronegativity down the group results in
 partially ionic bonding. This is not strong in SiC and whilst it tends to
 give co-ordination numbers of 4, can nonetheless result in several allotropes
 of decreasing symmetry: 4c, 4h, 6h. However, as the difference in period
 increases for the as yet theoretical
 group IV compounds, so too does the difference in electronegativity and hence
 also the bond ionicity and the degree of order. For a hot carrier
 absorber this is ideal because it is just such a large difference in the
 period which is needed to give the large mass difference and hence large
 phononic gaps. All of GeC, SnSi, SnC (and the Pb compounds) have computed
 phononic gaps large enough to block Klemens decay, and should also tend to
 form ordered diamond structure compounds.

 
	
  

 	
  

 
	
 b)

 	
 Because
 of their covalent bonding, the group IV elements have relatively small
 electronic band gaps as compared to their more ionic III-V and much more
 ionic II-VI analogues in the same period: e.g. Sn 0.15eV, InSb 0.4eV, CdTe
 1.5eV. In fact to achieve approximately the same electronic band gap one must
 go down one period from group IV to III-V and down another period from III-V
 to II-VI: e.g. Si 1.1eV, GaAs 1.45eV, CdTe 1.5eV. This means that for group IV compounds there is greater scope for large
 mass difference compounds whilst still maintaining small electronic
 band gaps. A small band gap of course being important for broadband
 absorption in an absorber - property 3 in the desirable properties for hot
 carrier absorbers listed above.

 
	
  

 	
  

 
	
 c)

 	
 The smaller Eg would tend to be for the
 larger mass compounds of Pb or Sn. Which, to give large mass difference, would be compounded with
 Si or Ge. This trend towards the lower periods of group IV also means that
 the maximum optical phonon and maximum acoustic phonon energies will be
 smaller for a given mass ratio - the desirable properties 4 and 5.

 
	
  

 	
  

 
	
 d)

 	
 Furthermore,
 unlike most groups, the group IV elements remain abundant for the higher mass
 number elements – desirable property 9. Property 10 is also satisfied because
 the group IVs have low toxicity.

 


Nanostructures:


As
discussed in Section 4.5.3.3.9, the phonon dispersion of QD nanostructures can
be calculated in the same way as compounds.
Their phononic properties can be estimated from consideration of their combination
force constants. Hence it is possible to ‘engineer’ phononic properties in a
wider range of nanostructure combinations.
Of the materials discussed above the Group IVs lend themselves most readily
to formation of nanostructures instead of compounds due to their predominantly
covalent bonding, which allows variation in the coordination number. Therefore
the nanostructure approaches of Section
4.5.3.3.9 are consistent with a similar description as analogues of InN,
whether it be III-V QDs, colloidally dispersed QDs or for core shell
QDs.


4.5.3.3.5 Combined phononic / mnw absorber design


Researchers: Gavin Conibeer, Yu Feng, Hongze Xia, Robert Patterson, Suntrana Smyth


The
multiple nano-well structures have demonstrated reduced carrier cooling by enhancing
the phonon bottleneck, probably at a stage prior to optical phonon decay. Large
phononic band gap materials would seem to have strong potential to block
Klemens’ decay of optical phonons. A structure combining both structures should
give even greater reduction in carrier cooling rates as the mechanisms should not interfere directly with
each other and hence their effects should be additive. A combined
absorber for a hot carrier cell should look similar to Fig. 4.5.33.


This
has narrow nano-wells of phononic band gap material, not thin enough to give
quantum confinement, separated by thin barriers. The barriers would either have
an electronic band offset to block hot carrier diffusion or an optical phonon
energy offset with the nano-wells and probably a coherent nano-well structure,
or all three, depending which of the multiple nano-well (MNW)


- 202 -






mechanisms
discussed above is dominant. This will enhance the phonon bottleneck either by
preventing hot carrier diffusion, by reflecting and confining optical phonons
or by blocking certain folded phonon modes, respectively, or all three. The
phononic nano-wells themselves will further enhance phonon bottleneck by
preventing Klemens’ decay of optical to acoustic phonons. This should maximise
the phonon bottleneck and minimise carrier cooling for a given illumination
intensity. A lattice matched pair of large phononic band gap nano-well and
barrier materials (InN/InGaN) is suggested in Fig. 4.5.33 (b), whereas in (c) a
wider range of possible InN analogue phononic band gap materials could be used
in a thin film structure with thin probably oxide barriers. The thin barriers
(a few nm) should facilitate tunnelling between wells and hence transport of
the carriers to the contacts.





Figure 4.5.33: Hot Carrier cell design combining Multiple
Nano-Well (MNW) with phononic band gap materials: (a) schematic and (b) band
structure for InN/InGaN epitaxial multiple nanowell with phononic band nano-wells; (c) band structure with
nano-wels of analogues of InN with large phononic band gap with thin
film barriers.


4.5.3.3.6 Alternative quantum well structure with integrated
energy selective contacts


Researchers: Dirk KÖnig, Binesh Puthen-Veettil


Use of quantum wells in the absorber region has the
advantage of a continuous density of states in the plane above the first confined energy level.
This allows continuous absorption of photon energies above this energy. With
barriers of different phonon energy such that they reflect optical phonons, phonons can be confined in the z direction and
increase phonon bottleneck the consequent decrease in carrier cooling
rate. If in addition the materials and barrier are chosen such that they give
two confined energy levels below the barrier a significant distance apart, then
these can select only a very few carrier energies to be transported through the
device, thus effectively playing the role of energy elective contacts. This
thus reduces the need for external energy selective contacts and allows the
contacts to be optimised for the appropriate Fermi level and work function
alignment to give electron collection and hole reflection at one contact and
hole collection and electron reflection at the other, thus giving the
rectification required.


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Figure 4.5.34: InAs/AlSb MQW on AlSb with GaP top electron
contact layer to give integrated hot carrier absorber and carrier energy
selected structure without need for external ESCs.


Figure
4.5.34 shows such a structure with InAs wells and AlSb barriers, designed to
give confined energy levels such that the 2nd
level aligns with the conduction band of heavily n-type GaP. The GaP and
AlSb contact materials act as hole and electron blocking layers respectively.
The valence band offset between InAs and AlSb is very small. This allows the
AlSb to collect thermalised holes from the valence band of the superlattice.
InAs and AlSb have a very small lattice mismatch (1.3%) so the whole structure
could be grown epitaxially on an AlSb substrate [ 4.5.59, 4.5.60].


Such
an integrated structure would give an ideal test bed for integration of
absorber MQW and internal energy selection of carriers which would not require
external ESCs and which gives rectification in the external circuit. The
specific materials and their thicknesses required would not allow optimisation
for slowing of hot carrier cooling, but would be able to be optimised for
appropriately spaced energy selection and collection.


4.5.3.3.7 Phononic dispersion modelling to predict phonon
lifetimes and decay rates in specific III-V or core-shell structures


Researchers: Hongze Xia, Yu Feng, Robert Patterson, Gavin Conibeer


InN/
InxGa1-xN multiple quantum-well superlattices (MQW-SL)
with wurtzite crystal structure are studied as the absorber of the hot carrier
solar cell. Such a structure will exploit both the significant
hot-phonon-bottleneck effects of the component materials and the known slowed
carrier cooling in MQWs. The former is due to the large contrast of the atomic
masses, and hence a large band-gap between high-lying and low-lying phonon modes.
The phonon band-gap effectively stops Klemens decay, i.e. one high-lying phonon
decaying into two low-lying phonons, and produce hot population on the polar
modes, feeding energy back to electrons. InN and InxGa1-xN
have very similar lattice structures, with
almost the same lattice constant. This benefits the solar energy conversion, in
terms of both the carrier transport and the reduction of recombination
sites. The 1-dimensional superlattice structure ensures a continuous electronic
energy spectrum and hence a broad-band absorption. The absorption is further
enhanced by the small electronic band-gap of InN.


The dispersion relations of phonon modes in
superlattices have been computed with a 1-dimensional atomic-plane model. Since the epitaxial
growth is intended to be along the high-symmetrical Γ - A direction for
both InN and InxGa1-xN layers, the periodicity of MQW-SL
and hence the zone-folding is along this
direction. For simplicity we assumed that the mode frequency of each
zone-folded mini-band only depends on the vertical component of the
wave-vector (i.e. along Γ - A). The reason for this assumption partly
comes from the high concentration of electron-emitted polar phonons around the
zone-center, where the dispersions are relatively flat compared to the mini-gaps.
In wurtzite


- 204 -






structure periodic atomic planes with alternating
elements align perpendicular to the Γ-A direction. If only considering the vertical component of
wave-vectors, all atoms in one atomic plane vibrate in phase and can be treated
as one uniform displacement. The phononic model adopted here treats a
1-dimensional chain with an equal length of the superlattice periodicity,
taking into account the plane-to-plane force constants. The plane-to-plane
force constants are calculated by adopting the conventional Keating potentials
for all bonds. A sample phononic dispersion relation is shown in Fig. 4.5.35.
The electronic structures, involving the wavefuntions and energies of all the
states, can be calculated by using the Kronig-Penney model for superlattices.
In equilibrium conditions space charges occur in the well layers (negative) and
in the barrier layers (positive) due to the difference of their electron affinities. The electron affinity
of InGaN is significantly lower than that of InN, giving a larger
off-set of the conduction band edge rather than that of the valence band edge.
A sample electronic dispersion relation is shown in Fig. 4.5.36.





Figure 4.5.35: The left figure indicates the phonon decay
paths and the sample dispersion for a SL structure with 3 layers of unit cells
inside each layer (barrier or well). The right figure shows the phonon
modulation function of a representative mode from each category of modes,
computed from the eigenfunctions of the lattice dynamic equation. The transverse
modes have similar modulation functions and hence only that part of the
longitude modes is shown.





Figure 4.5.36: The electronic dispersion relation and the
zone-center wave functions for a SL structure with 3 layers of unit cells
inside each layer (barrier or well). The wave vector is along the Gamma-A
crystal direction.


The calculation of the rate of polar interaction between
hot electrons (here emission by holes is not of concern) and polar phonons are based on the
Frohlich-type Hamiltonian and the 1-st order


- 205 -






perturbation theory. A hot reservoir of electrons at
1000K is assumed, with heat transferred to the cold reservoir of lattice modes
at 300K. The energy relaxation times referring to the polar phonon emission are illustrated in Fig. 4.5.37, for all the
MQW-SL configurations, i.e. for different well and barrier thicknesses and
different barrier materials. 18X18 combinations of the thicknesses are sampled
for representing all the possible structures from 2nm to 9nm. The number of
combinations comes from the fact that a
complete well/barrier layer should include an integer number of unit cells.
Figure 4.5.37 is generated by interpolating the relaxation time data of
the sampling combinations.





Figure 4.5.37: Energy relaxation times of the hot electron
system (1000K), in the reservoir of phonons
(300K), with superlattice structures of different well/barrier thicknesses and
InxGa1-xN compositions (Left: x=0 Right: x=0.2)


The
energy relaxation times of high-lying longitude optical phonons are
demonstrated in Fig. 4.5.38, for different combinations of well and barrier
thicknesses. For Indium mole fraction x=0 (left figure), the contrast of
relaxation times are larger than that for the case for x=0.2 (right figure);
this is reasonable for as the higher Indium content in the barrier layer would
make the structure closer to bulk InN. Both
figures show the same trend with different barrier/well thicknesses: the
relaxation time increases with thicker well layers and thinner barrier layers,
in spite of some irregular local variations. The regular variation
mainly results from the change of numbers of InN-like modes and InGaN-like
modes. Here a hot reservoir of high-lying modes at 1000K is assumed, with heat
transferred to the cold reservoir of low-lying lattice modes at 300K. The
non-equilibrium between the two systems is physical due to their significant
frequency separation.





Figure 4.5.38: Energy relaxation times of the high-lying LO
phonon system (1000K), in the reservoir of 1000K high-lying phonons and 300K
low-lying phonons, with superlattice structures of different well/barrier
thicknesses and InxGa1-xN compositions (Left: x=0 Right:
x=0.2)


To
explain the variation, we need to first examine the phonon dispersions of the
MQW-SL. Taking x=0 as an example, the left figure in Fig. 4.5.35 shows the
dispersions of InN/GaN SL structure, with


- 206 -






6
layers of nitrogen atoms inside each layer (the same for both the InN layer and
the GaN layer). From the number of atomic layers involved we could expect the
same number of high-lying optical modes which are InN-like and GaN-like
respectively. The GaN-like optical modes (blue color: solid line for LO, dash
line for TO) have much higher frequencies than the InN-like optical modes
(green color), for the bond force constants
of GaN is larger than for InN and the atomic mass of Ga is smaller than
In. There are also four branches (orange colour: two of LO and two of TO)
corresponding to the interfacial modes (IF), with frequencies between GaN-like
modes and InN-like modes. From the right figure
of Fig. (4), the vibrations can be seen to be almost completely confined in the
respective layers, for InN-like and GaN-like optical modes. For IF most
of the vibrational energy is within the 1st and 2nd nitrogen atoms from the interface, and hence has little overlap
with InN-like or GaN-like modes.


Considering
all three-phonon processes, only the Ridley channel and the Shrivistava-Barman
channel (decaying into a high- lying optical phonon and a low-lying optical
phonon) are allowed due to the energy conservation law. As in
wurtzite-structured SL, the difference between the acoustic branches and the low-lying optical branches becomes almost
indistinguishable with both having partial optical-like character. These
low-lying branches can be separated into two categories: GaN-confined modes (red lines) and GaN-InN mixed modes (grey lines).
In fact the InN-like low-lying modes sit within the allowed frequency
band of GaN, hence vibrations in GaN layers are excited too. Therefore no InN-confined modes exist in the low-lying branches.
Among all allowed three-phonon processes, GaN-like optical modes can
decay into both types of low- lying modes (See the solid arrow and the dashed
arrow in Fig. 4.5.35), while InN-like optical modes can only decay into the
mixed modes as they only overlap with the mixed modes. Besides, the energy gap
between GaN-like LO modes and GaN-like TO modes is relatively large; hence the
resulting low-lying phonons have relatively high energies, compared to the decayed
low-lying phonons from the InN-like modes. Since the low-lying branches with
high energies are generally flatter, leading to a larger joint density of
states of transition, the decay rates of GaN-like LO modes could be enhanced
further. Due to the two reasons explained above the GaN-like LO modes decay
faster than the InN-like LO modes. Therefore with a thicker well layer (or a thinner barrier layer), the
energy relaxation time of the high-lying LO phonon system becomes
longer, for it introduces more InN-like modes (or fewer InGaN-like modes).


According
to Fig. 4.5.38, the relaxation time of high-lying phonons could go up to 300ps,
significantly longer than the bulk value of around 1ps. And according to Fig.
4.5.37 the relaxation time of hot electrons
(if phonons are completely thermalized) can go up to more than 1ps, this is
also much longer than the bulk value. Combining these two, an optimized
MQW-SL structure should involve a thin barrier layer and a thick well layer.
The contrast in phonon energies between the well and the barrier also needs to
be large, indicating a small Indium content in the barrier. In addition, such a
structures could potentially prevent hot phonons diffusing out, making it even
more attractive than bulk materials.


4.5.3.3.8 Time resolved photoluminescence measurements of
bulk phononic band gap materials


The
potential efficiency boost, which can be achieved by hot carrier solar cells,
is directly related to the possibility of extracting high energy carriers from
the absorber layer before thermalisation, increasing
the voltage and hence the conversion efficiency. The poor conversion efficiency
of photons with energies above the band gap of the absorber is the main
loss mechanism in conventional single junction
solar cells. The investigation of thermalisation time constants of hot carriers
is a crucial step towards the engineering of hot carrier cells. The
efficiency of an InN based hot carrier solar cell has been calculated using the
theoretical model in Sections 4.5.3.3.7 and 4.5.3.1. It was found that the
limiting efficiency is strongly related to hot carriers relaxation velocity in
the absorber [4.5.61].


A
comparison of femtosecond time resolved photoluminescence (tr-PL) spectroscopy
between InP and GaAs was reported in the annual report two years ago and now
published in [4.5.48]. This showed a
distinctly longer carrier cooling time constant for the wide phononic gap InP
as compared to almost zero phononic gap of GaAs. It also showed further
evidence for excitation into higher side valleys for both GaAs and InP for
appropriate excitation wavelengths. Hot carrier cooling in InN has


- 207 -






been
investigated using tr-PL. The wide gap between optical and acoustic branches in
the InN phononic dispersion relation (wider than that for InP) prevents the
Klemens decay of optical phonon into acoustic phonons. This can lead to slower
carrier cooling due to the “Hot Phonon Effect” [4.5.62]. The decay of hot
carriers for different excitation wavelengths InN has been investigated.





Figure 4.5.39: Schematic representation of time resolved
photoluminescence arrangement.


Tr-PL
experiments have been performed on an InN sample grown on sapphire substrate,
obtained from collaborators in Taiwan, using the measurement configuration
shown in Fig. 4.5.39. In this technique a laser pulse acts as a switching gate
relating the photoluminescence signal to the time domain. The PL signal is
collected from the sample, after a femtosecond laser excitation, and focused in
a non linear crystal. The gate signal is generated from the same laser and is
focused on the same crystal after passing through an optical delay stage. The
signal is detected using a monochromator and a
PMT. Our system configuration provides 150 fs pulses with tunable wavelength
over a range of 256 nm (4.84 eV) to 2.6 μm (0.48 eV).





Figure 4.5.40: 3D representation of InN time resolved PL
data.


Figure
4.5.40 is a three dimensional representation of PL as a function of time for
all the probed wavelengths. It can be observed that the PL sharply rises when
the carriers are photo-excited by the laser pulse. The fast decay of the PL
shows the thermalisation of carriers towards respective band


- 208 -






edges.
The decay is faster for highly energetic carriers compared to carriers closer
to the bandgap. Thus the carrier population quickly degenerates towards the
band edges during the thermalisation process. In InN the thermalisation is most
probably due to interaction of highly energetic electrons and holes with LO
phonons.


To
investigate the rate of the carrier cooling process, the effective temperature
of the carrier population has been calculated fitting the high energy tail of
the PL spectrum for every single time during the cooling transient. The PL has
been fitted assuming that carriers form a Boltzmann-like distribution in a
femtosecond time scale using the following equation.





L(ε)
represents the measured PL intensity at energy ε, α(ε) is the
measured sample absorption coefficient, EG
is the InN energy gap, 0.7 eV in this case, and kB is the Boltzmann
constant. TC is the fitted parameter and represents the hot carrier
temperature. Figure 4.5.41 shows the tr-PL data for two sets of samples
(a) In N grown on Sapphire substrate and (b) InN grown by MBE at Saitama
University for zinc blende structure on sapphire and as wurtzite structure on
MgO substrate. For the data in (a) the
carrier temperature transient is seen to follow an exponential behaviour quite
well (blue line – fit). The data for the higher quality samples in (b)
show significantly slower carrier cooling rates than for (a).





Figure 4.5.41: Carrier relaxation curves for InN using
tr-PL: (a) Lower quality InN with soPhotoexcited carrier density is 1.5 x 1019
cm-3. The blue dashed curve is a single exponential fit. (b) Zinc
blende and wurtzite InN samples grown by MBE on MgO substrates by Shuei Yagi,
Saitama University (Cubic=crosses, Wurtzite structure = circles). This high
quality material shows much slower carrier cooling rates.


The
fitting of the calculated temperature data has been performed using a single
exponential.





Here
τTH represents the carriers thermalisation time constant,
whereas C and K are two constant parameters. The fitted value for τTH
is 7 ps. This long cooling constant can be attributed to the hot phonon effect
due to the long lifetime of the A1(LO) phonon [4.5.48,4.5.63]. The variation in
carrier cooling rates measured is attributed to the difficulty of growing good
quality InN on sapphire substrate, such that the carrier cooling velocity is
strictly related to the quality of the material. Other slow carrier cooling
constants have been reported for InN in the literature [4.5.64].


- 209 -






4.5.3.3.9 Nanostructures for absorbers


Nanostructures offer the possibility of modification of
the phonon dispersion of a composite material. III-V compounds or indeed most of the cubic
and hexagonal compounds can be considered as very fine nanostructures
consisting of ‘quantum dots’ of only one atom (say In) in a matrix (say N) with
only one atom separating each ‘QD’ and arranged in two interpenetrating fcc
lattices. Modelling of the 1D phonon dispersion in this way gives a close
agreement with the phonon dispersion for zinc-blende InN extracted from real
measured data for wurtzite material.


Similar
‘phonon band gaps’ should appear in good quality nanostructure superlattices,
through coherent Bragg reflection of modes such that gaps in the superlattice
dispersion open up [4.5.65]. There is a close analogy with photonic structures in which modulation of
the refractive index in a
periodic system opens up gaps of disallowed photon
energies. Here modulation of the ease with which phonons are
transmitted (the acoustic impedance)
opens up gaps of disallowed phonon energies.


Force constant modelling of III-V QD materials by SK growth


3D
force constant modelling, using the reasonable assumption of simple harmonic
motion of atoms in a matrix around their rest or lowest energy position,
reveals such phononic gaps [4.5.66]. The model calculates longitudinal and
transverse modes and can be used to calculate dispersions in a variety of
symmetry directions and for different combinations of QD sub-lattice structure
and super-lattice structure.


III-V
Stranski-Krastinov grown QD arrays of InAs in InGaAs and InGaAlAs matrices are
fabricated at the University of Tokyo using
MBE. We are investigating these for evidence of phonon dispersion modulation
and potential slowed carrier cooling. In order to understand the expected
phonon dispersions these are being modelled using the 3D force constant
technique.


Lattice
matched and strain compensated material pairs that may produce large phonon
bandgaps are of interest. Previous
iterations in the design of these structures indicated the importance of
separating “light” and “heavy” atoms to different parts of the
nanostructure. Initially, the lightest atom in the system, As, was present in
both the QD and the matrix. This meant that the reduced mass of both regions
(proportional to the sum of the inverses of each atomic mass) was very similar
as the light element dominates in this case.


On this iteration structures with an In0.5Ga0.5-xAlxAs
matrix (with x=0.4) and InAs QDs were grown. Significant Al content
was introduced into these structures with the expectation that this light
element, segregated to the matrix material, might produce appreciable phonon
bandgaps. Some images derived from
characterisation of the structure are presented in Fig. 4.5.42. The
superlattice of QDs has a simple hexagonal structure. Extraordinary periodic
out-of-plane stacking is achievable and largely defect free structures can be
grown on the order of microns.


- 210 -









Figure 4.5.42: Some images of the structure as grown and
modelled. (Above) the near-perfect stacking of the structure is shown. (Above right) in-plane InAs QD arrangement. (Right)
simple hexagonal superlattice structure showing stacking. The QDs are similar
to flattened disks with in plane dimensions
of 40-50 nm and heights of 7 nm.


Force
constant modelling of this structure predicts an appreciable phonon bandgap, as
shown in Fig. 4.5.43 and Fig. 4.5.44. This bandgap is due almost entirely to
the presence of the Al in the system. A small
bandgap is present due to the mass difference between In, Ga and As, but it is
less than a quarter of the size shown in Fig. 4.5.43 and Fig. 4.5.44.
Due to computational constraints the size of the QDs is quite small, only about
a nanometre in diameter. While actual sizes for these structures are too
computationally intensive to model without extreme effort, recent modelling
with gradually increasing size suggests that the dispersion relations scale
linearly with size. That is, once the discrete distances (bond lengths) are
small relative to the superlattice unit cell dimension, the dispersion should
look exactly the same when scaled such that the relative dimensions are
preserved. This will be investigated in greater detail in further work.





- 211 -









Figure 4.5.43: Dispersion curves for InAs QDs in a simple
hexagonal SL. The matrix is In0.5Ga0.5-xAlxAs
with (x = 0.4). There is a large phonon bandgap present in the system
due to the presence of Al. The gap seems very tolerant to Al location. (Top):
<100>, (Middle): <110>, (Bottom): <111> directions.





Figure 4.5.44: Densities of states (DOS) for the dispersion
curve immediately to the left. Note small changes in the DOS could be due to
pseudo-random locations of the Ga and Al particles or to differences in
crystalline direction. (Top): <100>, (Middle): <110>, (Bottom):
<111> directions respectively.


4.5.3.3.10 Fabrication and characterisation of highly ordered
nanoparticle arrays for hot carrier absorber


As demonstrated by the work on modelling phonon
modulation, periodic core-shell QD arrays offer a way to significantly change the phonon
modulation in a superlattice because the core and shell can be of materials of
very different force constant, directly leading to a strong phonon confinement.
Deposition methods are being investigated to fabricate such highly ordered QD
arrays.


The
Langmuir-Blodgett (LB) deposition technique is conventionally used for the
fabrication and characterisation of single
and multilayer films with precise control of thickness, molecular orientation and
packing density. In recent years, it has been applied to metallic nanoparticle
depositions for formation of highly ordered array.


We
have employed an LB system to fabricate gold (Au) and silicon (Si) nanoparticle
layers. A LB system has been set up at UNSW. We have designed the trough top
for nanoparticle deposition to optimise the dispersability. The system is shown
in Fig. 4.5.45, which consists of a trough, a pair of barriers, a surface
pressure sensor, a dipping unit for film transferring and an interface
controller.


- 212 -









Figure 4.5.45: A customised LB system
for nanoparticle deposition, installed in Room 1013 Chemical Engineering Building UNSW.


LB
deposition and formation of monolayer with Si nanoparticles:


Figure
4.5.46(a) is a schematic diagram showing the LB system, consisting of an LB
trough, a solid substrate, a dipping unit, a Wilhelmy plate and two barriers.
To make a Si NP array, a small amount of 0.5 mg/mL chloroform solution with
dodecane-capped Si NPs dissolved with diameters of 2.7 nm or 6.9 nm, was spread
onto the water surface drop wise over a time interval of 30 seconds. After the
complete evaporation of the solvent (chloroform), the Langmuir monolayer of Si
NPs on the water surface was dynamically
manipulated by compression with the barriers. During the compression, the
surface pressure (SP) was measured using the Wilhelmy plate, and the isotherm
curve (SP versus film area) was recorded. Fig. 4.5.46(b) shows the
isotherm curve for a 30 mL chloroform solution of 2.7-nm-diameter Si NPs with a
concentration of 0.5 mg/ml. Upon application of an appropriate level of
compression, suitable packing within the monolayer is achieved on the water
surface. Subsequently, the monolayer was transferred onto a hydrophilic
substrate (Si or Quartz wafer washed in piranha solution) through
vertical-dipping controlled by the dipping unit. From the AFM image (Fig.
4.5.46(c)), a film is observed. The cross-sectional view clearly shows the
thickness of the film is 2.7 nm, equal to the size of Si NPs, which indicates
a monolayer.





Figure 4.5.46: (a) a schematic of the LB system; (b)
measured isotherm curve for Si NPs of 2.7 nm and the chosen depositing surface
pressure (SP); (c) AFM image and cross-sectional profile of the fabricated 2.7
monolayer on a Si wafer.


- 213 -






Amorphous and short-range ordered close-packed arrays:


Unlike
other semi-conductor NPs such as CdSe or metallic NPs, due to their irregular
shapes and lighter weight, Si NPs are not mobile enough for self-assembly to
happen in a short timescale; this is due to the weaker attractive capillary
forces between the particles. Therefore, the reported results always show a
random array of Si NPs or aggregated clusters. In our work, we found that by
increasing the stabilizing time for monolayer formation on the water surface or
slowing down the moving rate of the barriers before performing the vertical
dipping, we can achieve decently uniform ordered close-packed arrays as shown
in Fig. 4.5.47 (d).





Figure 4.5.47: TEM images of (a) amorphous (b)
short-range-ordered (c) more uniform short-range-ordered close packed arrays
deposited at a SP of 50 mN/m with different LB processes (d) an enlarged area
in (c); Process 1 is to prolong the stabilizing time of the monolayer on the
water surface, 2 is to slow down the moving rate of the barriers.


In
summary, we have established a process for deposition of nanoparticle
monolayers on a solid substrate, with reasonably close packed ordered arrays of
nanoparticles. Future work will concentrate on improving periodicity of the
monolayer for both HC absorbers and energy selective contacts.


4.5.3.4 Implementation of nanostructures for hot carrier
cells


In
order to use a nanoparticle array as a hot carrier absorber, and hence utilise
slowed carrier cooling due to modified phonon dispersion, then a way to
fabricate a complete structure with ESCs must be established. A structure in
which nanoparticles are arranged in a uniform 3D array should give the required
phononic band gap. The degree to which this needs to be ordered is still under
investigation. Certainly a large difference between masses is beneficial
between the nanoparticles and the matrix


- 214 -






and
this can be best achieved in an array of ‘core shell QDs’ in which the core and
shell have very different acoustic impedance in order to promote coherent
reflection and hence confinement of phonons [4.5.65]. The matrix in which such
an array is embedded needs to allow transport of carriers to contacts and also
electron-electron and preferably hole-hole scattering to renormalise carrier
energies. Such a structure should also have a narrow electronic band gap so as
to absorb a wide range of photon energies.
This combination of properties is challenging but not mutually exclusive
because phononic properties are largely independent of electronic
properties and such a structure would not need
to have electronic confinement of carrier energies just vibronic modulation of
phonon energies. Energy selective contacts are also required for such a
structure. These would most ideally be arranged
at the top and bottom of the absorber. However, in a realistic device a
selective hole contact might not be required as the hole population only
contains a small fraction of the hot carrier energy and hence thermalisation of
holes is less important, whereas a structure with only one contact (presumably
the top contact) as an ESC would be much easier to fabricate. These ESCs would
be QD or QW double barrier structures. (Addition of extra layers to give double
QD/QW triple barrier structures (or even triple QD/QW quadruple barriers) would
give greater resonance and enebale the contact to be rectifying – important if it
is to collect carriers of only one. Such a QD array / double barrier QD ESC
structure is shown schematically in Fig. 4.5.48.





Figure 4.5.48: Conceptual integrated hot carrier device,
incorporating double (or triple) barrier QD resonant tunnelling contacts and
absorber probably based on either a bulk or nanostructure phononic band gap
material and/or MQW nanostructure, for slowed carrier cooling.


4.5.3.5 Summary of hot carrier solar cell research


There has been significant development in most areas of
hot carrier solar cell work. The modelling of efficiencies now includes combinations of
absorber and energy selective contacts with reasonable values chosen using real
material parameters based on III-nitrides. This has been carried out both for
bulk absorber InN and for multiple quantum well materials again based on
nitrides. Such devices are capable of being fabricated and then tested against
these predicted results and work with collaborators is moving towards growing
such devices. In addition the theory of the slowed carrier cooling behaviour in
MQWs is being investigated with plans to test the competing mechanisms again
using material from a range of collaborators. Measurement of InN with
time-resolved PL has indicated some evidence for slowed carrier cooling,
further corroborating the importance of a large phonon band gap to block
optical phonon decay, but also highlights the importance of material quality. Progress on the Langmuir-Blodgett approach to
ordered nanoparticle arrays has seen development of highly ordered arrays of
nanoparticles. The potential application of nanostructures to fully integrated devices
has been investigated conceptually, with various designs considered. Similarly
the possibility of absorber materials which are analogous to InN is also being
investigated. These many aspects of Hot carrier cells will see further
development and consolidation in 2013 with working devices planned for early
2014.


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4.5.4 Up-Conversion


Researchers: Sanghun Woo, Craig Johnson, Supriya Pillai, Shujuan Huang, Richard
Corkish, Gavin Conibeer


Collaboration with: Peter Reece (Physics, UNSW), John Stride (Chemistry, UNSW)


Up-conversion in novel silicon-based materials


Up-conversion
(UC) in erbium-doped phosphor compounds (particularly NaYF4:Er) has
been shown to be a promising means of enhancing the sub-band-gap spectral
response of conventional Si solar cells without modification of the electrical
properties of the cell [4.5.67]. In this scheme, a layer containing the
phosphor is applied to the rear of a high-efficiency bifacial cell. After
absorbing two long-wavelength (1500nm)
photons - which are transmitted by the cell - the excited Er ions can relax by
emitting a photon with an energy greater than the Si band gap, thereby
increasing the current that can be extracted from the cell.


Synthesis of Er doped NaYF4
phospors for UC


Fabrication of in-house phosphors allows control over
the growth process and the ability to control Er doping levels and the possibility to include
co-dopants.


NaY(1-x)F4:Erx
nanocrystals were synthesized by thermal decomposition of metal
trifluoroacetate (TFA) precursors [4.5.68, 4.5.69]. After the reaction at 800C
for two hours, oleic acid (OA) and 1-octadecene (1-ODE) were added to the
reaction and slowly heated for 1 hour at 2800C. At this stage, NaYF4:Er
NCs formed in the alpha phase. To convert NCs from alpha to the beta phase, the
temperature was slowly increased to 3200C for 3 hours. The resulting solution
was treated to remove unreacted precursors, excess OA and 1-ODE. NaYF4:Er
NC films were fabricated by spin casting followed by nucleophilic ligand
exchange. NC films were coated on the glass by multiple spin castings.


Fig.
4.5.49(A) shows HRTEM images of hexagonal beta phase NaYF4:Er NCs.
The dominant and most distinguishable lattice (100) of the hexagonal beta phase
NaYF4 NCs had a fringe distance of 0.54 nm. The histogram in Fig.
4.5.49 (B)indicates that particles had sizes ranging from 25 to 35 nm with a
narrow size distribution peaking around 30 nm.





Figure 4.5.49: Electron microscope
data for NaYF4: 10% Er. (A) low and high resolution TEM. Both scale bars correspond to 50nm. (B) Histogram
of diameters of Er dope NCs.


- 216 -






The absorption spectrum of UC NCs is shown in Fig.
4.5.50 and corresponds to all energy levels up to 3.1 eV of NaYF4:Er NCs. Fig.
4.5.50(B) shows five red-shifted peaks compared to the absorbance which were resulted from radiative relaxations
from 2H11/2, 4S3/2, 4F9/2,
4I9/2 and 4I11/2. As shown by UC luminescence
images in insets of Fig. 4.5.50(B) increased Er concentrations in UC NCs
accelerated multi-phonon relaxation processes between 4S3/2
and 4F9/2, which emphasizes the importance of
optimization between slowing non-radiative decay rates and increasing
absorption sites in NCs.





Figure 4.5.50: Absorption spectra for NaYF^4:Er NCs (A).
(B) detil 300-1000nm, with insets showing the luminescence.


The
SEM images of spin cast films before and after the soaking process in a MPA
solution are shown in Fig. 4.5.51(A) and (B) respectively. Without the post
soaking process, NCs were coated separately from neighbouring crystals, and
packed closely. By treating with a 3-MPA solution, NCs were agglomerated and linked to each other resulting
in a bulk-like feature. The cross-linked NC film made with 3-MPA brings
about easy handling, light trapping in NC layers, and a reduced film thickness.





Figure 4.5.51: (A) SEM images of morphology of upconverting
layer (A) before and (B) after soaking in a MPA solution. (C) Up-convesrion
spectroscopy showing the intensities of emitted photons in terms of laser power densities and Er doping concentration.
(D) Film thickness dependent upconversion under three different
excitation powers.


Figure 4.5.51 (C) shows the total intensity of
upconverted photons from energy levels above 4I11/2, which are all able to contribute to the
photoconversion process in c-Si solar cells, depending on the excitation power
between 30 μW and 215 mW and the Er concentration between 2 and 100% substituting Y sites of the NaYF4 host
structure. To minimize the effect of the total number of Er ions


- 217 -






in
UC samples, all samples were prepared thicker than the required thickness for
saturated upconversion, as described in the thickness dependent analysis below.
Considering the result of Er concentration effects, 15% Er doped UC layers
always gave the best upconversion regardless of the excitation power. This
means that increasing Er doping concentration did not affect the number of photons that could meet with each trivalent Er
ion. The ratio between the Er doping concentration and the photon flux
was not a predominant factor because the number of photons that could be
absorbed by each Er ion does not change when the films were thick enough. Thus,
increased or decreased Er content under fixed laser density did not have the
same effect as changing the laser power at the fixed Er concentrations. By
increasing Er densities in UC NCs, the number of neighbouring ions is increased, and the distance between Er ions is
reduced. Both can boost non-radiative cross relaxation processes as well
as energy transfer upconversion mechanisms. From the results, it can be concluded
that 15% Er doped NaYF4 can maximize the rate of the energy transfer
upconversion minus cross relaxation decay. This optimisation is practical
because it is independent of the solar concentrations.


Using the brightest UC NCs with a 15% doping concentration
of Er, the film thickness was optimized under various excitation fluxes, as illustrated in Fig. 4.5.51 (D).
From the excitation coefficient information,
the UC films had to be thicker than 20 μm, so that all photons within
energies of 4I13/2 could be absorbed by the UC
phosphor. However, the required film thicknesses to perform the saturated
emission were only about 2.5 μm under a 1 mW laser, 3 μm under a 5 mW
laser and 5 μm under a 10 mW laser. This is because upper UC NC layers
re-absorbed upconverted photons emitted from the lower layers. Also, the stable
bulk-like layers may have a higher absorption coefficient than simply packed
NCs, resulting in relatively thinner thicknesses required for the saturated UC
efficiency.


Light trapping in slowed light modes to enhance UC efficiency


While
phosphors have demonstrated high-efficiency UC behaviour, their use presents
particular challenges with regard to fabrication and cost. Last year we
presented data on use of Er-doped porous Si (PSi:Er) as an alternative UC
material. PSi is unique in that its porosity - and hence the material
refractive index - can be varied as a function of depth, allowing for the
elaboration of high-quality monolithic Si optical structures such as
distributed Bragg reflectors (DBRs).


A
cross-sectional electron micrograph of such a structure is shown in Fig.
4.5.52. Its porous substructure also allows for deep infusion of dopant atoms
via techniques such as electroplating.





Figure 4.5.52: Cross-sectional electron micrograph of a
30-bilayer porous silicon distributed Bragg reflector.


Figure
4.5.53 shows the calculated field intensity profile for these DBRs in the
spectral region between 1500 and 1600nm. The results have been normalized to
the incident field intensity and each


- 218 -






plot
is independently colour-scaled to show maximum detail. They clearly demonstrate
that considerable field enhancement is possible throughout the depth of such a
multilayer structure, with peak relative intensities of more than 11 for the
30-bilayer structure and more than 18 for the 40-bilayer structure. However, it is also clear that the regions of
enhancement are increasingly narrowly-concentrated as the number of bilayers
increases. This is due to the additional interference fringes in the
reflectivity characteristic of the DBR that arise with additional layers, as is
apparent from Fig. 4.5.53. Calculations for more than 40 layers show that
enhancements of up to 80 are possible for a large number of bilayers though the
peak narrows to less than 1nm wide for 100 bilayers.





Figure 4.5.53: Intensity-valued mode profile in DBRs with
varying numbers of high/low-porosity bilayers
for a leading band edge fixed at 1550nm. All intensities are normalized to the
incident field intensity and each plot is independently color-scaled to
show maximum detail. The air and substrate interfaces with the DBR are
indicated by the dashed black (top) and red (bottom) lines, respectively.


Sensitisation of Er phosphors to a wider wavelength range


Er has an absorption window at the I13/2
level from 1480nm to 1580nm. This narrow range means that it is inefficient at absorbing below Si band
gap photons. Sensitizing to the 1100-1500nm range would significantly enhance
the absorption and hence have a dramatic effect on the up-conversion
efficiency, because of its non-linear dependence.


PbS
nanoparticles have a wide absorption range below 1100nm. They can also be tuned
in size to emit at 1500nm, with absorbed photons being downshifted to 1500nm at
some quantum efficiency less than 100%.


Such PbS NCs have been synthesized based on the
procedure developed by Cademartiri et al. [4.5.70] by decomposition of PbCl2 with
sulphur solution. PbS QDs and Er-doped NaYF4 NC films were


- 219 -







fabricated
by the simple spin casting method followed by nucleophilic cross-linking of
NCs.


The
sizes, the size distribution and the fringe distance of resulting UC NCs were
determined using HRTEM and glancing incident x-ray diffraction (GIXRD) was
carried out to investigate the crystalline structures of PbS QDs.





Figure 4.5.54: Absorbance and emission data from UV-Vis
spectroscopy. (a) for PBs quantm dots showing the desired slight mismatch
between absorbance and emission such that re-absorption is unlikely. (b)
absorbnce for Er doped NCs and emission from PbS QDs, showing the good match
between PbS emission to pump Er absorption at the.


Figure 4.5.54 shows absorbance and emission data for PbS
QDs and Er doped NaYF4 NCs. The range of absorbance for PbS QDs is
close to ideal for absorbance of light below the Si bandgap but stopping short of the emission. This emission, tuned
to 1500nm, is ideal for pumping the I13/2 level at 1520nm in Er,
without being re-absorbed by the PbS. Thus the sensitization of Er to a
significantly wider wavelength range should be very valuable in increasing the
flux of photons which can be absorbed and up-converted by the Er doped
phosphor.


Conclusion


Up-conversion
in Er doped phosphors shows very promising results. Fabrication of NaYF4 NCs allows greater control over doping levels and
geometries for efficient up-conversion. Light trapping in DBR structures
based on P-Si should enable greater absorbance of 1500nm light through the
greater density of photonic states. Progress towards sensitisation of Er to a
wider wavelength range using down shifting PbS quantum dots tuned to emit at
1500nm has been made. Combination of all these elements in an up-conversion structure is expected to significantly
boost UC efficiencies for Si cells.


4.5.5 Concluding Remarks for the Third Generation Section


Work
has proceeded significantly in all the areas of Third Generation research, with
improved fabrication and characterisation of materials and complexity of
modelling which together give an overall better understanding and optimisation
of devices.


Group
IV based nanostructure materials have seen significant improvement in
understanding of the parameters required for growth and the influence of these
on material and device properties. The modelling of equivalent circuits for the
devices is giving greater understanding and now allowing prediction of growth
conditions to give improvement. The realisation that the effective bandgap and hence the VOC in these materials is
dominated by defects has led to a greater effort on characterising and
reducing defects. Light trapping through back reflectors and the use of
plasmonics has led to more than double the current in nanostructure devices.
Further optimisation in this area is very likely. The use of conductive
substrates to reduce the lateral resistivity has made promising progress, with
SiC and Mo substrates showing some improvements. Optimisation of these will
continue to give better lateral transport and should lead to better photovoltaic
properties. Higher VOCs over 500mV


- 220 -






have
been measured at lower temperatures. The justification for this is that due to
the high resistivities, devices heat up considerably and cooling is valid to
bring them back to room temperature. A temperature of 50degress below ambient
is calculated to achieve this condition and gives
VOC of 503mV, although significantly high VOCs occur at
even lower temperatures as would be expected. Work on other related
materials has seen lower resistivities for Si3N4
interlayers in Si QDs in SiO2; high quality Ge QW growth; initial
data for growth of Ge:C materials with the aim of engineering band gaps through alloying and the potential for transfer
of growth processes to PECVD with it lower levels of damage and thermal
budget.


Hot
carrier cells have seen a great deal of advance in the last year. Modeling of
overall devices now includes a much larger range of real material properties in
a more integrated structure. In addition modelling of combined multiple quantum
well an phononic bandgap structures is leading to new designs for combined structures which offer promise for slowed carrier
cooling. Modelling of energy selective contacts has progressed again
with real material properties and integration into full conceptual devices.
Such devices are entirely feasible and plans to grow these with collaborators
are underway. New designs for rectifying ESCs or incorporation of ESCs within
the absorbed material have been suggested and these also can be incorporated
into real structures. Increased understanding of the mechanisms involved in
both phononic band gap materials and in multiple quantum well structures is
leading to new experiments to determine the exact relationship between reduced
carrier cooling mechanisms. The results of these will lead to the optimum
design requirements for combined structures which should allow significantly
slower carrier cooling rates. Deposition of nanoparticle arrays has made progress with very ordered arrays
now possible using Langmuir-Bodgett deposition. These arrays have been modelled
to give significant phonon bandgaps and characterisation of these is now
underway. New materials suitable for some of these phononic and nanostructure
properties are being identified and growth and characterisation of these
materials is planned in-house and with a range of collaborators.


Up-conversion
has seen a further significant development with synthesis of NaYF4 doped with
Er now allowing greater flexibility in doping and geometry of devices. Light
trapping work is continuing to increase the
number of photonic modes available for absorption in Er. Sensitisation of Er to
a wider wavelength range is being enabled using PbS QDs which are now being
synthesised and characterised and which are showing close to ideal
properties for such sensitisation. Combination of all these approaches is
expected to give significantly improved up-conversion.


The
development of all the 3rd generation projects now allows much
greater understanding of the materials and devices. Work in 2013 will see
consolidation of this into improved devices. Several new areas of funding will
contribute to this and also allow development of new project areas.


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 Guillemoles, D. KÖnig, S. Shrestha, R. Clady, M. Tayebjee, T. Schmidt, M.A. Green “Hot Carrier solar cell
 Absorbers”, 24th European PVSEC, Hamburg (Sept. 2009).

 
	
  

 	
  

 
	
 4.5.47

 	
 F.
 Chen, A.N. Cartwright, Appl Phys Lett, 83 (2003) 4984.

 
	
  

 	
  

 
	
 4.5.48

 	
 R.
 Clady, M.J.Y. Tayebjee, P. Aliberti, D. Konig, N.J. Ekins-Daukes, G.J.
 Conibeer, T.W. Schmidt, M.A. Green, submitted PIP, 20 (Jan 2012) 82–92. DOI:
 10.1002/pip.1121.

 
	
  

 	
  

 
	
 4.5.49

 	
 G.
 Conibeer, R. Patterson, P. Aliberti, H. Xia, S. Huang, K. KÖnig, B.
 Puthen-Veettil, S. Shrestha, M.A. Green, “Hot carrier cell absorber
 materials,” Proc. 26th European Photovoltaic Solar Energy Conference Hamburg,
 (2011).

 
	
  

 	
  

 
	
 4.5.50

 	
 Y.
 Rosenwaks, M. Hanna, D. Levi, D. Szmyd, R. Ahrenkiel, A. Nozik, “Hot-carrier
 cooling in GaAs: quantum wells versus bulk,” Phys Rev B, 48, 14675-14678
 (1993).

 
	
  

 	
  

 
	
 4.5.51

 	
 J-F.
 Guillemoles, G. Conibeer, M.A. Green, “On achievable efficiencies of Hot
 carrier solar cell absorbers based on reported experimental thermalisation
 properties,” Proc. 21st European Photovoltaic Solar Energy Conference,
 Dresden Germany, 234–237 (2006).

 
	
  

 	
  

 
	
 4.5.52

 	
 L.
 Hirst, M. FÜhrer, A. LeBris, J-F. Guillemoles, M. Tayebjee, R. Clady, T.
 Schmidt, Y. Wang, M. Sugiyama, N.
 Ekins-Daukes, Proc. 37th IEEE PV Specialists Conference, Seattle, talk
 927 (2011).

 
	
  

 	
  

 
	
 4.5.53

 	
 C.
 Colvard, T.A. Gant, M.V. Klein, “Folded acoustic and quantized optic phonons
 in (GaAl)As superlattices,” Physical Review B, 31, 2080-2091 (1985).

 
	
  

 	
  

 
	
 4.5.54

 	
 T.K.
 Todorov, K.B. Reuter, D.B. Mitzi, Advanced Materials, 22 (21010) 1–4.

 
	
  

 	
  

 
	
 4.5.55

 	
 W.R. L. Lambrecht, “Structure and phonons of ZnGeN2”,
 Phys. Rev. B 72 (2005) 155202

 
	
  

 	
  

 
	
 4.5.56

 	
 T.R.
 Paudel, W.R.L. Lambrecht, Phys. Rev. B, 79 (2009) 245205.

 
	
  

 	
  

 
	
 4.5.57

 	
 S.C.
 Erwin, I. Zutic, Nature Materials, 3 (2004) 410-414.

 
	
  

 	
  

 
	
 4.5.58

 	
 Z.T.
 Liu, J.Z. Zhu, N.K. Xu, X.L. Zheng, Jap. J. Appl. Phys., 36 (1997) 3625.

 


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 4.5.59

 	
 D.
 Konig, Y. Takeda, B. Puthen-Veettil, G. Conibeer, Japanese Journal of Applied
 Physics 51 (2012) 10ND02, Lattice-Matched Hot Carrier Solar Cell with Energy
 Selectivity Integrated into Hot Carrier Absorber.

 
	
  

 	
  

 
	
 4.5.60

 	
 D.
 Konig, Y. Takeda, B. Puthen-Veettil, Appl. Phys. Lett., 101, 153901 (2012),
 Technology-compatible hot carrier solar cell with energy selective hot
 carrier absorber and carrier-selective contacts.

 
	
  

 	
  

 
	
 4.5.61

 	
 P. Aliberti, Y. Feng, Y. Takeda, S.K. Shrestha, M.A.
 Green, G.J. Conibeer, “Investigation of theoretical efficiency limit of hot carrier solar cells with bulk InN
 absorber”, Journal of Applied Physics, 108 (2010) 094507-10.

 
	
  

 	
  

 
	
 4.5.62

 	
 K.T.
 Tsen, J.G. Kiang, D.K. Ferry, H. Lu, W.J. Schaff, H.W. Lin, S. Gwo, “Direct
 measurements of the lifetimes of longitudinal optical phonon modes and their
 dynamics in InN”, Applied Physics Letters, 90 (2007) 3.

 
	
  

 	
  

 
	
 4.5.63

 	
 Y.C.
 Wen, Y.C., C.Y. Chen, C.H. Shen, S. Gwo, C.K. Sun, “Ultrafast carrier
 thermalization in InN”, Applied Physics Letters, 89 (2006) 3; M.D. Yang, et
 al., “Hot carrier photoluminescence in InN epilayers”, Applied Physics A -
 Materials Science & Processing, 90 (2008) 123-127.

 
	
  

 	
  

 
	
 4.5.64

 	
 F. Chen, A.N. Cartwright, H. Lu, W.J. Schaff,
 “Ultrafast carrier dynamics in InN epilayers”, Journal of Crystal Growth, 269 (2004)
 10-14.

 
	
  

 	
  

 
	
 4.5.65

 	
 G.
 Conibeer, N.J. Ekins-Daukes, J-F. Guillemoles, D. KÖnig, E-C. Cho, C-W.
 Jiang, S. Shrestha and M.A. Green, “Progress on Hot Carrier Cells”, Solar
 Energy Materials and Solar Cells, 2009, 93, 713.

 
	
  

 	
  

 
	
 4.5.66

 	
 R.
 Patterson, M. Kirkengen, B. Puthen Veettil, D. Konig, M.A. Green, G.
 Conibeer, Sol. Ener. Mats. and Sol. Cells, 94 (2010) 1931-1935.

 
	
  

 	
  

 
	
 4.5.67

 	
 A.
 Shalav, B.S. Richards, M.A. Green, “Luminescent layers for enhanced silicon
 solar cell performance: Up-conversion”, Sol. Energy Mats. and Solar Cells, 91
 (2007) 829-842.

 
	
  

 	
  

 
	
 4.5.68

 	
 J.-C.
 Boyer, L. A. Cuccia, and J. A. Capobianco, “Synthesis of Colloidal
 Upconverting NaYF4: Er3+/Yb3+ and Tm3+/Yb3+ Monodisperse Nanocrystals,” Nano
 Letters, vol. 7, pp. 847-852, 2007/03/01 2007.

 
	
  

 	
  

 
	
 4.5.69

 	
 C.
 RÜssel, “A pyrolytic route to fluoride glasses. I. Preparation and thermal
 decomposition of metal trifluoroacetates,” Journal of Non-Crystalline Solids,
 vol. 152, pp. 161-166, 1993.

 
	
  

 	
  

 
	
 4.5.70

 	
 Cademartiri, L., et
 al., Multigram Scale, Solventless, and Diffusion-Controlled Route to Highly
 Monodisperse PbS Nanocrystals. The Journal of Physical Chemistry B, 2005.
 110(2): p. 671-673; Cademartiri, L., et al., Size-Dependent Extinction
 Coefficients of PbS Quantum Dots. Journal
 of the American Chemical Society, 2006. 128(31): p. 10337-10346.

 


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