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Title

湿式化学成長した超極薄AlドープSiOxにより形成された
Al誘起酸化物反転層を有する新規Si系太陽電池

Author(s)中島, 寛記CitationIssue Date

2024-03

Type

Thesis or Dissertation

Text version

ETD

URL

http://hdl.handle.net/10119/19072

RightsDescription

Supervisor: 大平 圭介, 先端科学技術研究科, 博士

Japan Advanced Institute of Science and Technology

Doctoral Dissertation

Novel Si-based Solar Cells withan Al-induced Charged Oxide Inversion Layer
Formed by a Wet Chemically Grown
Ultrathin Al-doped SiOx

Hiroki Nakajima

Supervisor: Keisuke OhdairaJapan Advanced Institute of Science and Technology
Division of Advanced Science and Technology
Materials Science

March 2024

Referee in Chief:Professor Keisuke OhdairaJapan Advanced Institute of Science and Technology

Referees:
Associate Professor Masashi Akabori
Japan Advanced Institute of Science and Technology
Professor Toshikazu Suzuki
Japan Advanced Institute of Science and Technology
Professor Hiroshi Mizuta
Japan Advanced Institute of Science and Technology
Dr. Hitoshi Sai
National Institute of Advanced Industrial Science and Technology

ABSTRACT
Carrier-selective passivating contacts would realize excellent surface passivation for crystalline Si (c-Si) solar cells witha high conductivity without direct metal electrode contacts, causing a deterioration of the interface properties. Hence, itwould be a breakthrough technology to access a practical limit of solar cell performance.
This thesis addresses using a simple process to realize low-cost and high-efficiency c-Si solar cells with asimple device structure of carrier-selective passivating contacts whose emitter is induced by external electrostatic charges.
The main major contribution of this work is developing a novel surface passivation method using a simple wet
chemical process, which does not require vacuum, high temperature, or hydrogen treatment processes, which are essential
processes in industrial c-Si solar cells. The process is just immersing Si wafers into a hot aluminum nitrate (Al(NO 3)3)
aqueous solution to form an ultrathin Al-doped SiOx layer possessing hole selectivity. On a planar n-type c-Si surface, themaximum effective surface recombination velocity (Seff, max) reached as low as 16 cm s−1. The analyzed saturation current
density (J0) of 65 fA cm−2 and contact resistivity (ρc) of 20 mΩ cm2, the hole selectivity (S10) value could be estimated to
be 13.3 which is competitive to other selective contacts created using vacuum and high-temperature process. Since theprocess eliminates the complexity of solar cell fabrication processes with an excellent carrier-selective passivating contact
property, this study suggests that a wet chemical process using an aqueous solution of Al(NO3)3 can be an industrial solar
cell fabrication process. In addition, the mechanism of the hole-selective passivating contact property was revealed. The
field-effect passivation is mainly attributed to the surface passivation.
The second contribution of this work is a demonstration of the device operation of a novel device structurewith an inversion layer induced by the negative fixed charges existing in ultrathin Al-doped SiOx. This result implies thatthe charge-induced inversion layer acts as a p+ emitter formed by a simple wet chemical process. Hence, using a surface
charge-induced emitter, this type of solar cell realizes high-throughput, low-cost p+ emitter formation by a simple wet
chemical process. This novel solar cell was named an Al-induced charged oxide inversion layer (Al-COIL) solar cell. The
conversion efficiency (η) of >10% was achieved for the device fabricated on a pyramidal textured Si substrate. In
particular, an excellent FF of 0.794 was realized using a fine Ag grid electrode on the emitter side. The reduction due toa high emitter sheet resistance could be mitigated by improving carrier collection through a fine Ag grid electrode covering
a relatively larger emitter surface area. However, still low Jsc and Voc were confirmed. PC1D simulation implies that the
reductions might originate from carrier recombination at the emitter surface due to the large number of interface states
and the increased minority carrier (electron) density by the diminished fixed charge density (Qf) less than −1×1011 cm−2.
The third contribution of this work is a discovery that surface passivation, and the sheet resistance of ultrathin
Al-doped SiOx/Si improve by the light illumination, which is the variable for addressing the origin of the hole-selective
passivating contact property. The significant increment of these properties implies that the excited electron trapping in
the Al-induced acceptor states existing in the Al-doped SiOx layer forms negatively charged AlO4− structures and enhances
field effect passivation induced by the fixed charges.
Lastly, achieving an excellent surface passivation quality on a pyramidal textured c-Si surface was crucial. The
lower Seff, max was confirmed as low as 13.3 cm s−1, corresponding to iVoc of 682 mV and J0 of 59 fA cm−2, respectively.
The light-trapping enhancement and slight space charge region compression compensate for the relatively enormous
amount of interface state density due to the surface area increment on the textured Si surface, reducing surface
recombination velocity and increasing solar cell performance. Hence, this study provides insight into the novel surface
passivation mechanism on a practical surface structure for the c-Si solar cell. It requires a suitable surface structure design
with a nano-scale size for increasing a hole-selective passivating contact by effectively exploiting electron trapping andspace charge region compression.
Keywords: c-Si solar cells, carrier-selective passivating contacts, electrostatic charges, ultrathin Al-doped SiOx, wetchemical oxidationACKNOWLEDGMENT
I am immensely grateful to my supervisor, Professor Keisuke Ohdaira, for his invaluable guidance and support during my
Ph.D. research. His thoughtful guidance and generous financial and mental assistance enabled me to focus and do mybest for this research work. Without his help, I would not have been able to complete my Ph.D. thesis successfully.
I would like to express my gratitude to Professor Yoshifumi Oshima, my vice-supervisor, for advising me on myresearch proposal. I appreciate our productive discussions, which helped me achieve my research goals.
I would like to express my heartfelt gratitude to Professor Shinya Maenosono for his supervision of my sub-theme
research on Al-doped h-WO3 nanoparticles. I truly appreciate his enthusiastic guidance and valuable advice on myresearch policy and approach while conducting this study.
I am grateful to Assistant Professor Mari Takahashi. Her careful guidance was valuable for particle synthesis, TEM,
and XRD measurements.
I would like to express my gratitude to Associate Professor Masashi Akabori for his guidance and support in the clean
room of CNMT in JAIST. I am also thankful for taking time out of his busy schedule to review my thesis.
I would also like to acknowledge Professor Toshikazu Suzuki for taking time out of his busy schedule to serve as areviewer for this thesis.
I would also like to acknowledge Professor Hiroshi Mizuta for taking time out of his busy schedule to serve as areviewer for this thesis.
I would also like to acknowledge Dr. Hitoshi Sai from the Global Zero Emission Research Center, AIST, for taking
time out of his busy schedule to serve as an external reviewer for this thesis.
I would like to express my gratitude to Dr. Tatsuya Murakami from CNMT, JAIST, for his assistance in performing
XPS measurements and obtaining valuable data. Additionally, my sincere appreciation goes to Dr. Koichi Higashiminefrom CNMT, JAIST, for conducting TEM-EDS measurements and providing crucial data for this research.
I express my gratitude to Dr. Ayako Okano from Technical Service Office for Pure and Applied Science, University ofTsukuba, for her assistance in performing SPM measurements to discuss the surface passivation mechanism from thesurface potential on a pyramidal textured Si surface.
I also express my gratitude to Dr. Katsuto Tanahashi and Dr. Tomihisa Tachibana from the Fukushima Renewable
Energy Institute, AIST, for their assistance in evaluating the surface passivation quality of my samples using PV-2000A.
Moreover, I acknowledge Dr. Takuya Matsui from the Global Zero Emission Research Center, AIST, for his valuableadvice during our discussion on this study.
I am extremely grateful for the opportunity to conduct research at the Ohdaira laboratory, where I received valuableassistance from its members. I express my sincere appreciation to Assistant Professor Huynh Thi Cam Tu for her guidanceand support throughout my research. I would also like to extend my gratitude to the researchers, Mr. Shinichiro Hayashi
and Ms. Tomoko Hori, for their assistance in my research activities. I am also thankful for the help extended by thestudents in the Ohdaira laboratory, Mr. Zheng Wang, Mr. Shuntaro Shimpo, Mr. Ryota Ohashi, Mr. Pengyu Yang, Mr.
Yuhi Ito, Mr. Mitsunori Nagahara, Mr. Chenxi Li, and Mr. Baitong Li. I extend my gratitude to the following Ohdaira lab
graduates: Dr. Yuli Wen, Mr. Ryoichi Ishikawa, Mr. Wenbo Cao, Mr. Dequin Wu, Mr. Yujia Liu, Mr. Toshiki Itasaka, Mr.
Yudai Yanagi, and Mr. Kuan Liu. I express my gratitude to former research assistant Ms. Miyuki Katagiri in Ohdaira lab.
I acknowledge my previous supervisors at SANKEN, Osaka University: Professor Hikaru Kobayashi, Professor Masao
Takahashi, Associate Professor Taketoshi Matsumoto, and Associate Professor Kentaro Imamura. During my master's
program, I worked on the topic of surface passivation by nitric acid oxidation and gained extensive knowledge andexperience in wet chemical processes under their guidance. This led me to the concept for my current research. I would
also like to express my appreciation to Professor Kazuyuki Edamoto at Rikkyo University, whose instruction on XPS and
AES was beneficial for my research. Without their guidance, I could not have developed this research program.
This Ph.D. study was supported by JST SPRING, Grant Number JPMJSP2102. I would like to acknowledge theirfinancial support.

TABLE OF CONTENTS
1. Introduction……………………………………………………………………..………..1
1.1 Conventional crystalline Si solar cells and the recent trend…………………………….2
1.2 Promising solution for the cost-effective crystalline Si solar cells……………………...3
1.3 Potential of Al-doped SiOx material…………………………………………………….5
1.4 Outline of this work…………………………………………………………………….6
References……………………………………………………………………………………7

2. Surface recombination theory and passivating contacts……………………11
2.1 Carrier recombination in crystalline silicon1………………………………………….12
2.1.1

Carrier recombination in bulk………………………………………………....122.1.2

Carrier recombination at the surface…………………………………………..15

2.2 Carrier-selective passivating contact………………………………………………….20
2.2.1

Surface passivation…………………………………………………………….20

2.2.2

Theory of carrier-selective contacts……………………………………………21

2.2.3

Carrier selectivity……………………………………………………………...24

2.2.4

Passivating contact materials and structure…………………………………...26

2.3 Summary……………………………………………………………………………...28
References…………………………………………………………………………………..29

3. Characterization techniques.......................................................................................33
3.1 Carrier lifetime measurements......................................................................................33
3.1.1

Quasi-steady state photoconductance (QSSPC) ................................................33

3.1.2

Microwave photoconductivity decay (µ-PCD) ...................................................36

3.2 X-ray photoelectron spectroscopy (XPS) .....................................................................37
3.3 Current and voltage (I–V) measurement........................................................................40

3.4 Capacitance and voltage (C–V) measurement...............................................................44
3.5 Four-point probe technique...........................................................................................45
3.6 Fourier transform infrared spectroscopy (FTIR) ...........................................................47
3.7 UV-vis-NIR spectrometry..............................................................................................48
3.8 Scanning electron microscopy (SEM) ..........................................................................48
3.9 Transmission electron microscopy (TEM) ....................................................................49
3.10 Energy dispersive X-ray spectroscopy (EDS) ...............................................................49
3.11 Scanning probe microscopy (SPM) ..............................................................................49
References..............................................................................................................................51

4. Hole-selective contact by an ultrathin Al-doped silicon oxide......................53
4.1 Impurity doped SiOx for carrier-selective contacts........................................................54
4.2 Concept of aluminum nitrate oxidation treatment..........................................................54
4.3 Film properties of Al-doped SiOx..................................................................................56
4.3.1

Sample fabrication..............................................................................................56

4.3.2

Sample evaluation...............................................................................................56

4.3.3

Chemical bond states..........................................................................................57

4.3.4

Atomic concentration and thickness...................................................................58

4.3.5

Composition of sub-oxide and SiO2 component..................................................59

4.3.6

Al composition evaluated by surface sensitive condition....................................60

4.3.7

Atomic density of interfacial SiOx.......................................................................62

4.3.8

Cross-sectional structure....................................................................................64

4.4 Surface passivation quality by Al-doped SiOx...............................................................66
4.4.1

Sample fabrication..............................................................................................66

4.4.2

Sample evaluation...............................................................................................66

4.4.3

Surface passivation quality.................................................................................67

4.5 Electrical properties of Al-doped SiOx/Si interface.......................................................71
4.5.1

Tunneling current density...................................................................................71

4.5.2

Contact resistivity...............................................................................................74

4.5.3

Sheet resistance..................................................................................................75

4.5.4

Determination of Fixed charge density and Interface state density.....................77

4.6 Carrier selectivity of Al-doped SiOx..............................................................................80
4.7 Effect of surface orientation..........................................................................................81
4.7.1

Surface passivation on Si(100) and Si(111) ........................................................81

4.7.2

Chemical bond states on Si(100) and Si(111) .....................................................83

4.7.3

Electrical properties on Si(100) and Si(111) ......................................................85

4.7.4

Effect of light illumination..................................................................................87

4.8 Model of hole-selective passivating contact..................................................................90
4.9 Summary ......................................................................................................................91
References..............................................................................................................................91

5. Al-induced charged oxide inversion layer solar cells.......................................97
5.1 Concept of Al-COIL architecture..................................................................................98
5.2 Simulation for Al-COIL solar cells..............................................................................100
5.3 Fabrication and evaluation of Al-COIL solar cells.......................................................108
5.3.1

Sample fabrication............................................................................................108

5.3.2

Sample evaluation.............................................................................................109

5.3.3

J–V characteristics of Al-COIL solar cells........................................................110

5.4 Summary.....................................................................................................................113
References............................................................................................................................114

6. Anti-reflective structures and layers on Al-COIL architecture...................117
6.1 Strategy of improvement for Al-COIL solar cell performance.....................................118
6.1.1

Enhancement of photocurrent density...............................................................118

6.1.2

Enhancement of surface passivation quality.....................................................119

6.2 Effectiveness of pyramidal textured Si surfaces..........................................................122
6.3 Sample fabrication.......................................................................................................124
6.3.1

Texturing process..............................................................................................124

6.3.2

Cleaning and surface passivation.....................................................................124

6.4 Sample evaluation.......................................................................................................126
6.5 Surface morphology....................................................................................................128
6.6 Estimation of photon current density...........................................................................129
6.7 Surface passivation quality..........................................................................................131
6.8 Surface potential..........................................................................................................135
6.9 Fabrication and evaluation of solar cells......................................................................138
6.9.1

Sample fabrication............................................................................................138

6.9.2

Sample evaluation.............................................................................................139

6.9.3

J–V characteristics of Al-COIL solar cells........................................................139

6.10 Anti-reflective coating on Al-COIL structure..............................................................141
6.10.1 Sample fabrication............................................................................................142
6.10.2 Sample evaluation.............................................................................................143
6.10.3 Surface passivation of SiNy/Al-doped SiOx stacks............................................144

6.10.4 Fixed charge density of SiNy/Al-doped SiOx stacks...........................................145
6.10.5 Optical properties of Cat-CVD SiNy on pyramidal textured Si.........................149
6.11 Summary ....................................................................................................................150
References............................................................................................................................151
7.

Conclusion......................................................................................................................153
Appendix.........................................................................................................................155
Appendix 1: Simulation parameters for PC1D.....................................................................155
Appendix 2: Simulation parameters for PC3S......................................................................158
Appendix 3: Simulation parameters for OPAL 2..................................................................159
Appendix 4: Photon flux spectrum of solar radiation............................................................160
Appendix 5: Absorption coefficient of c-Si and penetration length in c-Si...........................161
Appendix 6: Inelastic mean free path (IMFP) ......................................................................162
Appendix 7: τeff measurement using low excitation laser light photon density......................163
Appendix 8: Surface passivation by different aluminum precursor......................................164
Appendix 9: Stability and recovery of surface passivation quality........................................165
Appendix 10: Surface passivation quality on p-type c-Si.....................................................167
Appendix 11: Frequency dependence of C–V measurements...............................................168
Appendix 12: τeff degradation by residual water....................................................................169
Appendix 13: Fabrication of inverted pyramidal texture......................................................171
Appendix 14: Surface passivation on inverted pyramidal texture.........................................174
Appendix 15: PEDOT:PSS on ultrathin Al-doped SiOx........................................................176
Appendix 16: Al/Mg/ultrathin Al-doped SiOx stacks............................................................178
Appendix 17: Surface passivation on rough Si surface.........................................................180
Appendix 18: PMMA on Al-doped SiOx for anti-reflection coating.....................................182
Appendix 19: WOz on Al-doped SiOx for anti-reflection coating..........................................185
Appendix 20: Fabrication of nano-sized pyramidal textured c-Si surface............................189
Appendix 21: Potassium nitrate (KNO3) treatment on c-Si surface......................................193
References............................................................................................................................195

List of publications.......................................................................................................196

CHAPTER 1

INTRODUCTION

CHAPTER 1
INTRODUCTION
It is commonly recognized that CO2 emissions are the primary factor of global warming [1],resulting in rising sea levels due to melting icebergs [2], intense storms like hurricanes [3],droughts causing significant crop injury [4], and the like. Since 1850, the cumulative CO 2
emissions reached 2400 ± 240 GtCO2 in 2019 [1]. In response, the Intergovernmental Panel on
Climate Change (IPCC) announced that 45% of global CO2 emissions must be reduced by the endof 2030 to keep the temperature increase below 1.5 °C [1]. Otherwise, an unchecked scenario will
await our future.
Most of the CO2 emission comes from coal, gas, and oil, which are energy sources [1]. Thus,
one of the promising solutions is using sustainable energy such as solar, wind, and waterpower
because these energy sources are inexhaustible and have extremely low CO2 emissions. In
particular, solar energy is a widely accessible energy source because the sunshine is evenly
distributed worldwide [5]. Researchers in Bell laboratories demonstrated the first practical silicon(Si) solar cell in 1954 [6]. Fortunately, solar cells are commercially available nowadays, and themarket is dominated by wafer-based crystalline Si (c-Si) solar cells with a 95% share [7]. Overseveral decades, the dramatic cost reduction has reached a comparable cost with fossil fuel,resulting in explosively rising photovoltaic (PV) installation and surpassing 1 TW as a cumulative
PV module shipment in 2022 [8]. However, further cost reduction and initiating more PV
installations will be essential to achieve our next goal of cutting 45% off greenhouse gas emissions
by the end of 2030 [9].

1CHAPTER 1

INTRODUCTION

1.1 Conventional crystalline Si solar cells and the recent trend
The several types of conventional c-Si solar cell structures are shown in Figure 1.1. The
mainstream c-Si solar cell technology is the Passivated Emitter and Rear Cell (PERC) emergedin 1989. This technology has allowed us to obtain high solar cell efficiency compared with AlBSF solar cells with a structure of fully Al-diffused and creating back surface field (BSF) at therear Si surface due to the rear design of surface passivation and local Al-diffused contact area [10].
According to the calculation by Institut für Solarenergieforschung in Hameln (ISFH), the PERC
cell efficiency limit will be 24.5% [11]. Since the theoretical c-Si solar cell efficiency limit is29.43% [12], 4.93% can be improved by other novel cell architectures using state-of-the-art
technologies. The following upcoming technologies that can be mainstream are expected to betunnel oxide passivating contact (TOPCon) [13] or Si heterojunction (SHJ) [14] solar cells due tothe higher theoretical efficiency limit of 28.7% and 27.5%, respectively [11]. Such highperformance can be realized by minimizing surface carrier recombination, which reduces directcontact between metal electrodes and Si surface while selecting carrier types at the surface using
field effect passivation, called passivating contact [15]. The 1.2% difference is mainly derived
from parasitic absorption loss in the amorphous Si (a-Si) front surface passivation layer used in
the SHJ structure [16]. In TOPCon structure, high bandgap materials with excellent opticaltransparency, such as aluminum oxide (Al2O3) and Si nitride (SiNx), can be applied as a surface
passivation layer on boron (B) diffused p+ front emitter, realizing very low parasitic absorption
loss of < 0.1 mA cm−2 on the front surface [17]. Recently, the Chinese PV manufacturer LONGi
recorded cell efficiency of 26.81% using SHJ technology [18], and Jolywood has accomplished26.7% using TOPCon technology [19]. Since TOPCon process cost is lesser than SHJ mainly due
to lower Ag consumption and the absence of a transparent conductive oxide (TCO) layer such as
indium tin oxide (ITO) for the TOPCon fabrication process, the market share of TOPCon is almost
double compared with SHJ, 15% in 2022 [20]. Moreover, TOPCon’s market share is expected to
be dominated by 60%, while SHJ will be 19% in 2033 [20].

Fig. 1.1 Device structures of Al-BSF, PERC, TOPCon, and SHJ solar cells.

2CHAPTER 1

INTRODUCTION

1.2 Promising solution for the cost-effective crystalline Si solar cells
In general, however, the B-diffused p+ emitter surface is challenging to be passivated, causing
relatively high energy loss of ~0.4 mW cm−2 at the p+ emitter [21], compared with ~0.1 mW cm−2
for SHJ [18]. Previous research showed that the B-diffusion process forms the boron-rich layer(BRL). Since thermal B-diffusion forms misfit dislocations and causes recombination centers [22,23], BRL would be lifetime killers. Thus, the BRL removal process is necessary, and several
additional processes have been developed, as reported elsewhere [22, 23]. Moreover, the energy
state level created by the boron oxygen (B-O) complex, located near the Si mid-gap [24], highly
degrades solar cell performance. In addition, much higher temperatures are required for the Bdiffusion process compared to the P-diffusion using POCl3 precursors due to the low diffusion
coefficient of B into Si substrates [25]. Therefore, the lower fabrication throughput and higher
energy consumption are recognized as a drawback of TOPCon technology compared withconventional PERC. Some breakthrough novel technology will be required to prevail overpassivating-contact c-Si solar cells manufactured by high-throughput processes and lower energy
consumption with a low-temperature route.
A recent study proposed the dopant-free asymmetric heterocontact (DASH) concept anddemonstrated high-efficiency c-Si solar cells of over 20% [26]. This type of solar cell possesses
transition metal oxide with a high work function, such as MoO3 [27], WO3 [28], and V2O5 [29],instead of a B-diffused layer. The upward band bending of the Si surface is induced by contacting
such high work function materials on an n-type c-Si surface, creating an inversion layer with alarge amount of hole (minority carrier) accumulations at the Si surface. However, the sharp
reduction of electrons by repelling from the Si surface allows us to obtain hole-selective contact[30]. On the other hand, low work function materials such as LiFx [31], TiOx [32], and MgOx [33]contacting on an n-type Si surface create an accumulation layer with a large number of electrons(majority carrier), but the sharp decrease in a hole concentration at the Si surface, enabling
electron-selective contact. Therefore, since the inversion layer plays a role in the p+ emitter of ntype c-Si solar cells without impurity diffusion into the Si surface, the impurity-rerated defectstates are not formed. Moreover, this DASH technology does not require the high-temperatureprocess [26].
In addition, an inversion layer can be created using an oxide layer like AlOx or Al-doped SiO2
possessing extremely high fixed charge density (Qf) on the order of 1012 cm−2 [34, 35]. For
instance, Al-doped ultrathin SiO2 has a negative Qf, inducing hole accumulation and creating an
inversion layer at the surface of n-type c-Si. This Al-doped ultrathin SiO2 layer on n-type c-Si
demonstrated excellent surface passivation quality with a saturation current density (J0) as low as
25 fA cm−2 corresponding to an effective surface recombination velocity (Seff) of 2.7 cm s−1 and
3CHAPTER 1

INTRODUCTION

contact resistivity (ρc) of ~200 mΩ cm2 [36]. Furthermore, the inversion layer induced by Al2O3
showed an excellent lateral hole conductivity on n-type c-Si, which effectively acts as an emitter
layer and demonstrated more than 20% efficient n-type c-Si solar cells [37, 38]. However, in thisprevious research, since atomic layer deposition (ALD) was adapted to form Al2O3, high vacuum
condition is required, and energy consumption and the maintenance cost of flammable gases liketrimethyl aluminum (TMA) will be much higher [39]. As an alternative, atmospheric pressurechemical vapor deposition (APCVD) is a high-throughput industrially comparable deposition
process [40]. However, the relatively high temperature of ~400 °C and >10 nm thick Al2O3 layer
is necessary to obtain an excellent surface passivation quality [40], which cannot be used for
carrier-selective passivating contacts due to such a thick dielectric layer.
A novel wet chemical technology that might be a breakthrough has emerged. E. Schmid et al.
presented a cost-effective negative fixed charge passivation method, the so-called “Wet AlOx”
process, in 2018 [41]. This process requires two wet chemical processes: formation of ultrathin
SiOx dipping cleaned Si wafers in hot sulfuric peroxide mixture followed by dipping in aqueous
solution dissolved Al salts (Na[Al13O4(OH)24(H2O)12](NO3)8, AlCl3, or Al(NO3)3). Through these
wet chemical processes, the layer structure with one monolayer Al adsorbed on wet-chemicallyformed SiOx was realized with a high negative surface charge of up to ~1.4×1012 cm−2. The selfsaturation of Al adsorption results in monolayer deposition of AlOx. In addition, they
demonstrated cell efficiency comparable to conventional silicon nitride (SiNx) passivated samples
on n-type c-Si with a higher tolerance to many metal impurities. Although such an ultrathin
negative fixed charge layer has a considerable potential to realize a hole-selective passivating
contact, a wet chemically grown ultrathin negative fixed charged oxide layer has not demonstrated
the device operation.
Therefore, this Ph.D. research demonstrated an outstanding surface passivation quality despite
an ultrathin oxide layer of ~1 nm thickness with excellent hole-selectivity by a novel, simple wet
chemical process, and the device operation was successfully confirmed with an induced inversion
layer acting as a p+ emitter of n-type c-Si solar cells [42, 43]. I named this novel solar cellarchitecture “Al-induced charged oxide inversion layer (Al-COIL) solar cells.” It allows
simultaneous p+ emitter formation and surface passivation in a simple low-temperature wet
chemical process. However, one of the drawbacks is that this process cannot form a relatively
thick layer for an antireflective coating (ARC). When this layer is used as a front surface of solar
cells, some layers with good optical properties, such as the SiNx layer, must be deposited, or an
ultralow reflection structure, such as black Si [44], formed to obtain higher solar cell performance.
However, due to plasma damage to the ultrathin layer, such an ultrathin layer is unlikely to deposit
ARC by conventional plasma-enhanced chemical vapor deposition (PECVD) [45]. Hence, theARC has to be gently deposited by other alternatives. As shown in Figure 1.2, a comparison of4CHAPTER 1

INTRODUCTION

process flows for PERC, TOPCon, SHJ, and interdigitated back contact (IBC) [46] with Al-COIL,
Al-COIL has a massive advantage with much fewer processes compared to any otherindustrialized process flows. However, it is noted that this Al-COIL process flow is provisional.
In this work, since I focused on the study of Al(NO3)3 treatment process, metal electrodes wereformed by physical vapor deposition (PVD) as a fundamental research step. Hence, it is necessary
to develop a rear-side surface passivation and screen printing (SP) process for the mass productionof Al-COIL solar cells in the future.

Fig. 1.2 Process flows for PERC, TOPCon, SHJ, IBC, and Al-COIL.

1.3 Potential of Al-doped SiOx materialFirstly, since Al, Si, and O are highly abundant elements on earth, and the top three elements on
Clarke number: 49.20% for O, 25.67% for Si, and 7.50% for Al, respectively [47], choosing Aldoped SiOx is a sustainable option. Secondly, a sufficient Al concentration of ~6×1014 cm−2 iscontained in one monolayer AlOx [48]. Suppose 100% of Al atoms have negative charges onthem ; in that case, an impeccable surface passivation quality will be expected from the calculationof Seff as a function of Qf, as shown in Figure 1.3. The calculation was performed using a simplified
Shockley–Read–Hall (SRH) equation for a single defect energy level under the assumption of
sufficiently high negative Qf [49] assuming the electron capture cross section (σn) of 1.8×10−14
cm2 [40] and interface state density per area (Nit) of ~1.0×1012 cm−2 [40] for resistivity of 2.5 Ω
cm and Δn of 1×1015 cm−3 condition. Even though the layer is one monolayer with only 6.7%
negatively charged Al, the calculated Seff will be the same value obtained from PECVD a5

CHAPTER 1

INTRODUCTION

Si/SiOx/SiNx stacks, which are reported as the most excellent surface passivation for c-Si surface
as Seff of 0.1 cm s−1 [50]. Furthermore, since four coordinated aluminum oxides (AlO4) adjacentto the SiO2 structure stabilize in the four coordinated SiO2 structures [48], creating a negative
fixed charge, the field effect passivation from Al-doped SiOx material is the highly expected
approach.

Fig. 1.3

Relationship between effective surface recombination velocity (Seff) and fixed charge

density (Qf) in Al-doped SiOx/Si structure.

1.4

Outline of this work

This doctoral dissertation accesses the cost-effective and high-efficiency novel Si-based solar cells,
realizing an excellent surface passivation quality and hole-selective contact by an ultrathin Al-doped
SiOx, which can form using a simple wet chemical treatment. Its structure is divided as follows.
Chapter 2 reviews the principle of carrier recombination in c-Si and its surface passivation with carrierselective contacts. Chapter 3 explains the characterization techniques this work used. Chapter 4
demonstrated excellent surface passivation quality and the hole-selective contact by a simple wet
chemical process. Moreover, several approaches, such as spectroscopic, electrical, and microscopic
methods, revealed the mechanism of the hole-selective passivating contact. Chapter 5 demonstrated
the operation of a solar cell device with Al-COIL architecture, the p-type inversion layer acting as ap+ emitter. Chapter 6 investigated and discussed the effectiveness of a pyramidal-textured Si surface
and an anti-reflection coating to increase cell efficiency.
6

CHAPTER 1

INTRODUCTION

References[1]

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INTRODUCTION

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E. Schmid, S. Schmitt, T. Boescke, E. Wefringhaus, F. Buchholz, C. Peter, R. Marczak, A.
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10

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTSCHAPTER 2
SURFACE RECOMBINATION THEORYAND PASSIVATING CONTACTS

Surface recombination is one of the crucial events in Si-based devices. Nowadays, since the bulk
property has been exceptionally high and the wafer is becoming thinner, the surface is much morecritical for the device's performance. Also, pyramidal textured and nano-structured Si surfaces areused to boost photocurrent density, but surface recombination will be a much more crucial
problem due to the increase of surface defect density derived from the larger surface area.
However, if the surface is well passivated by induction of Si band bending, the carriers can beeffectively separated, causing carrier-selective contact. This might be a breakthrough if a simple
process achieves it. Hence, it is essential to understand the surface recombination theory andconstruct the mitigative methods.
In this chapter, to provide a theoretical background for the c-Si solar cells with carrier-selective
contact, the recombination theory in the bulk and surface and the surface passivation method arereviewed. In particular, the description is focused on the carrier-selective contacts. Several typesof carrier-selective contacts are reviewed to understand this research finding of hole-selective
contact by a charged oxide. A simplified case of the surface recombination model, proposed by
McIntosh and Black, followed for this work, is described. This chapter also introduces a methodfor determining the Qf and the interface state density values for ultrathin dielectric layers throughthe combination of measurable values of surface potential and the surface recombination model.

11

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

2.1 Carrier recombination in crystalline silicon
2.1.1

Carrier recombination in bulk

When a semiconductor absorbs light with a more significant energy than the material's band gap
energy, it generates excited carriers. This process is called generation. However, the energy must
be released as photons or phonons to return the system to thermal equilibrium by combining withthe excited electrons and holes. These recombination processes can be measured as carrier
lifetime (τ), the time it takes for the carriers to recombine. τ can be described as

𝜏=

𝛥𝑛
.
𝑈

(2.1)

Where Δn is the excess carrier density, and U is the recombination rate per volume. The
recombination process takes place in the bulk and surface.

Radiative recombination
The radiative recombination directly occurs by releasing excess energy as a photon with an energy
corresponding to the close value of band gap energy, and the radiative recombination rate can be
described as𝑈rad = 𝐵rad (𝑛𝑝 − 𝑛i2 ).

(2.2)

Where Brad is the radiative recombination coefficient, n is the electron density, p is the hole density,and ni is the intrinsic carrier density. The Brad value of Si is known as 9.5×10−15 cm3 s−1 at room
temperature [1]. The value is considerably lower than the direct transition material, such as GaAs,
due to the indirect transition material of Si. Therefore, since the contribution of the radiative
recombination of the Si material is low compared with other types of recombination, it can be
negligible.

Auger recombination
Auger recombination is also known as intrinsic recombination. The recombination process isinvolved in a three-carrier interaction. The recombination occurs by a relaxation due to the energy
transfer of one excited electron (or hole) to the neighbor excited electron (or hole) to excite them12CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

higher state. The Auger recombination rate can be described as
𝑈Aug = 𝐶n (𝑛2 𝑝 − 𝑛i2 𝑛0 ) + 𝐶p (𝑛𝑝2 − 𝑝0 𝑛i2 ) ≈ 𝐶n 𝑛2 𝑝 + 𝐶p 𝑛𝑝2 .

(2.3)

Where Cn and Cp are the Auger coefficients for electron and hole. The Auger lifetime of p-typematerial can be simplified as

𝜏Aug,li =11
2 𝑎𝑛𝑑 𝜏Aug,hi = (𝐶 + 𝐶 )𝛥𝑛 2 .
𝐶p 𝑁A
n
p

(2.4)

These are Auger lifetimes at low and high injection levels, respectively. From these equations,Auger lifetime at a high injection level or heavy doping is crucial to the Si solar cells. Therefore,
the doping concentration in the emitter must be optimized so as not to decrease the Auger lifetime.
However, in the actual case, the Auger lifetime is lower than the generated value from Eq. (2.4)
due to the columbic interaction of charge carriers. One of the correct models was proposed by
Kerr and Cuevas [2]. The model created by fitting the experimental data set is given by

𝜏Aug =𝛥𝑛.
𝑛𝑝(1.8 × 10−24 𝑛00.65 + 6 × 10−25 𝑝00.65 + 3 × 10−27 𝛥𝑛0.8 )(2.5)

The more accurate parametrization was proposed by Richter et al., which is consistent with thetheory of Coulomb-enhanced Auger recombination and accounts for Coulomb-enhanced radiative
recombination [3].

Shockley-Read Hall (SRH) recombination
The Schockley-Read Hall recombination is the recombination via a defect state in the band gapof the material. The recombination process was first investigated by Schockley, Read, and Hall[4, 5]. Therefore, the recombination process is called Schockley-Read-Hall (SRH) recombination.
The SRH recombination is known as extrinsic recombination due to the recombination comingfrom extrinsic factors such as extended defects: Dislocations or stacking faults and intrinsic or
extrinsic point defects, causing the discrete energy levels in the band gap.
Firstly, an electron in the conduction band relaxes to the defect level and afterward relaxes tothe valence band, annihilating a hole. The defect state near the mid-gap of the material drasticallyincreases the recombination rate because of the similar probability for the capturing of electrons
from a conduction band and holes from the valence band, which is referred to as the recombination
13

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

center. The recombination rate based on the SRH model with a single defect level is given by

𝑈SRH =

𝑛𝑝 − 𝑛𝑖2,𝜏𝑝0 (𝑛 + 𝑛1 ) + 𝜏𝑛0 (𝑝 + 𝑝1 )𝑛1 = 𝑛i exp (𝜏𝑛0 =𝐸t −𝐸i),𝑘B 𝑇

𝑝1 = 𝑛i exp {1,
𝑁t 𝜎n 𝜈th𝜏𝑝0 =−(𝐸t −𝐸i )
},
𝑘B 𝑇

1
.
𝑁t 𝜎p 𝜈th

(2.6)

(2.7)

(2.8)

Where n1 and p1 are the statistical values of the equilibrium electron and hole density when thefermi level is consistent with the defect energy level (EF = Et). τn0 and τp0 are the capture time
constants for the electron and hole. The SRH recombination lifetime (τSRH) can be expressed as

𝜏SRH =

𝜏n0 (𝑝0 + 𝑝1 + 𝛥𝑛) + 𝜏p0 (𝑛0 + 𝑛1 + 𝛥𝑛).
𝑛0 + 𝑝0 + 𝛥𝑛

(2.9)

Where σn and σp are the capture cross section for electron and hole, respectively, and νth is thecarrier thermal velocity. At the low injection level, the τSRH is independent of the Δn and can be
simplified as, for p-type,𝜏SRH = 𝜏n0,(2.10)

𝜏SRH = 𝜏p0 .

(2.11)

for n-type,

At the high injection level, since the Δn is considerably higher than n0, n1, p0, and p1, τSRH is also
independent of the Δn and can be simplified as
𝜏SRH = 𝜏n0 + 𝜏p0 .

14(2.12)

CHAPTER 2

2.1.2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

Carrier recombination at the surface

The carrier recombination at the surface is different from bulk as follows: the active defect levels
at the surface are continuously distributed over the entire band gap. As an extended SRH
formalism, the surface recombination rate (Us) [6, 7] is described as
𝐸C
𝐷it (𝐸)2
𝑈S = (𝑛s 𝑝s − 𝑛i )𝜈th ∫𝑑𝐸.
𝐸V 𝑛s + 𝑛1 (𝐸) + 𝑝s + 𝑝1 (𝐸)
𝜎p (𝐸)𝜎n (𝐸)(2.13)

Where ns and ps are the electron and hole densities at the surface, and Dit(E) is the interface state
density per energy interval. By using Us, the surface recombination velocity (SRV) can be
described as

𝑆=

𝑈S
.
𝛥𝑛s

(2.14)

Where Δns is an excess carrier density at the surface, hence, the recombination lifetime at thesurface can be defined as an inverse SRV. The Eq. (2.13) is generally simplified by assuming asingle defect level with the interface state per unit area (Nit) as

𝑛s 𝑝s − 𝑛i2
𝑈S = 𝑛 + 𝑛
𝑝s + 𝑝1 ,
s
1
+
𝑆
𝑆
p0

(2.15)n0

𝑆n0 = 𝜎n 𝑁it 𝜈th and 𝑆p0 = 𝜎p 𝑁it 𝜈th .

(2.16)

Where Sn0 and Sp0 are the SRVs of electron and hole, respectively. The simplified SRH equation,
proposed by McIntosh and Black [8], was applied to this research. In the simplification to the ntype Si, the following situations are restricted: ns >> n1, (ns + n1)/Sp0 >> (ps + p1)/Sn0, which means
the defect is not very close to the band edges. The situation of psns >> 𝑛i2 is also imposed, which
is valid in all situations but deficient injection level. Hence, under these assumptions, the SRH
equation can be simplified as
𝑈S = 𝑆p0 𝑝S .

15

(2.17)

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

The equilibrium minority carrier density at the distance of d: the edge of the energy band,
unaffected by changes in the surface dielectric layer, is negligible. Hence, the carrier density at
the distance of d can be described as𝑛d = 𝑁dop + Δ𝑛d.

(2.18)

Where Ndop is the doping density in bulk, and Δnd is the excess carrier density at the distance ofd. Since the minority carrier is generated as an excess carrier, the Δnd can be replaced with aminority carrier density at d (nd),(2.19)

𝑛d = 𝑁dop + 𝑛d .
Therefore, by using Eq. (2.14) and (2.17), the effective SRV can be described as

𝑆eff = 𝑆p0

𝑝s
.
𝑝d

(2.20)

Thus, under the negligible band bending condition, ps should be equal to nd, and Seff is equal to
Sp0. When a positive charge is applied, ps < pd, leading to Seff < Sp0. Vice versa, when a negative
charge is applied, ps > pd, leading to Seff > Sp0.
The relationship between the carrier densities, band bending, and surface charge [6] in Eq. (2.21).

𝑝s − 𝑝d + 𝑛s − 𝑛d + (𝑁A − 𝑁D )𝛹s𝑄2=
𝑉T 2𝑞𝜀Si 𝑉T(2.21)

Where Ψs is the surface potential with respect to Ψd = 0, εSi is the permittivity of Si, and Q is the
charge in the material induced by the dielectric layer and the interface. The quasi-Fermi levels
between the surface and the position at a distance of d are assumed flat in Eq. (2.21).
The carrier densities can be described using the surface potential. When assuming Boltzmann
statics and the flat quasi-Fermi levels, surface carrier densities can be described as𝑛s = 𝑛d exp (

16

𝛹s
),
𝑉T

(2.22)

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

𝑝s = 𝑝d exp (−

𝛹s
).
𝑉T

(2.23)

Therefore, the relationship between surface carrier densities of surface and the distance of d canbe described as,(2.24)

𝑝s 𝑛s = 𝑝d 𝑛d .
For the exceptional cases below, Seff can be derived as follows,(1) In the case of negligibly small Qf (Ψs ~ 0),𝑛s = 𝑛 d ,(2.25)

𝑆eff = 𝑆p0 .

(2.26)(2) In the case of sufficiently small Qf, form Eq. (2.21) and (2.23), since n-type material, thisequation requires ns, nd >> ps, pd, simplifying Eq. (2.21) to give

exp (

𝛹s
𝑁D 𝛹s
𝑄2
)−1−
=
.
𝑉T
𝑛d 𝑉T 2𝑞𝜀Si 𝑉T 𝑛d(2.27)

In low injection level, when pd ~ NA, it can be further simplified by introducing a partial Maclaurin
series for carrier densities,
𝛹 k
(𝑉s )
𝛹s
𝑄2
∑ T −1+=,𝑘!
𝑉T 2𝑞𝜀Si 𝑉T 𝑁A∞(2.28)

k=2

𝛹s
𝑄
=
,
𝑉T √𝑞𝑉T 𝜀Si 𝑁A

𝑛s = 𝑛d exp (𝑄√𝑘B 𝑇𝜀Si 𝑁A) and 𝑝s = 𝑝d exp (−17(2.29)𝑄√𝑘B 𝑇𝜀Si 𝑁A),(2.30)

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

𝑆eff = 𝑆p0 exp (−𝑄√𝑘B 𝑇𝜀Si 𝑁D).

(2.31)(3) In the case of sufficiently positive Qf, carrier densities of Eq. (2.21) equals ns (ns >> nd, ps,pd), therefore,

𝑛s =

𝑄2
,
2𝑘B 𝑇𝜀Si

𝑆eff = 𝑆p0 𝑛d

2𝑘B 𝑇𝜀Si.
𝑄2(2.32)(2.33)(4) In the case of sufficiently negative Qf, carrier densities of Eq. (2.21) equals ps (ps >> ns, nd,pd), therefore,

𝑝s =

𝑄2
,
2𝑘B 𝑇𝜀Si

𝑆eff = 𝑆n0 𝑛d

2𝑘B 𝑇𝜀Si.
𝑄2(2.34)(2.35)

By substituting the Eq. (2.16),

𝑆eff = 𝜎n 𝑁it 𝜈th 𝑛d

2𝑘B 𝑇𝜀Si.
𝑄2(2.36)

Therefore, Nit can be described as

𝑁it =𝑆eff 𝑄 2.
2𝑘B 𝑇𝜀Si 𝜎n 𝜈th 𝑝d(2.37)

Thus, if Q and Seff are obtained, Nit can be determined using Eq. (2.37) since the other parameters
are known values.

18CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

This study used a highly negatively charged oxide layer on the n-type c-Si substrate. Hence, Eq.
(2.37) was adequate for the Al-doped SiOx/n-type Si substrate, and the equation was used todetermine the Nit value with analyzed Seff from carrier lifetime measurement. The σn was referredfrom the reported value in the literature [9–11]. The σn of the interface trap between a dielectric
layer of SiO2 or Al2O3 and Si have a similar value of ~10−14 cm2 at the mid-gap of Si. There is no
significant difference among these different processed samples, which has been reported inseveral literatures [9–11]. Hence, in this study, the value of 1.8×10−14 cm2 for Al2O3/p-type c-Si
[11] was applied to determine Nit. Therefore, Seff and Q must be experimentally obtained to accessthe Nit value. The SRV (S) can be described by introducing effective minority carrier lifetime (τeff)
as below.
1
1
1
1
1
=
+
+
+
,
𝜏eff 𝜏rad 𝜏Auger 𝜏SRH 𝜏surf
1
𝜏bulk

=

1
𝜏rad

+

1
𝜏surf

1
𝜏Auger

=

+

1
𝜏SRH

,

2𝑆
,
𝑊

(2.38)(2.39)(2.40)

1
1
2𝑆
=
+
.
𝜏eff 𝜏bulk 𝑊(2.41)

Where τbulk is the bulk lifetime, and W is the thickness of the substrate. For the n-typesemiconductor, the value of Sp0Δps can be defined as the recombination rate per unit area (cm−2s−1), and it can be described as a form of the diffusion equation for the Δn from surface to bulk ofmaterial when the carrier generation evenly occurs.

𝑆p0 𝛥𝑝(0) = 𝐷p

𝑑𝛥𝑝(𝑥)
𝑑𝑥𝑥=0(2.42)

Where Dp is the diffusion coefficient for hole. The τbulk has been modeled [2] as,for n-type,1𝜏bulk= (𝛥𝑝 + 𝑁D )(1.8 × 10−24 𝑁D 0.65 + 3.0 × 10−27 𝛥𝑝0.8 + 9.5 × 10−15 ), (2.43)
19

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

and for p-type,1𝜏bulk= (𝛥𝑛 + 𝑁A )(6.0 × 10−25 𝑁A 0.65 + 3.0 × 10−27 𝛥𝑛0.8 + 9.5 × 10−15 ) . (2.44)

Hence, the τbulk of the used Si substrates in this study can be estimated to be 12.6 ms with aresistivity of 2.5 Ω cm for n-type Si substrate at the Δn of 1×1015 cm−3 using Eq. (2.43). For thedetermination of Q value, from the Eqs. (2.23) and (2.34), Q is proportional to the root of
exponential Ψs as described,

𝑄 = √2𝑘B 𝑇𝜀Si 𝑝d exp (−

𝛹s
).
𝑉T

(2.45)The surface potential corresponds to surface band bending, which can be analyzed from
capacitance–voltage (C–V). The evaluation technique is described in the next Chapter.

2.2 Carrier-selective passivating contact
2.2.1 Surface passivation
Since the surface recombination centers are distributed in the entire band gap of the material, thesurface recombination significantly affects the solar cell performance, especially in open-circuit
voltage (Voc). To achieve high-efficiency solar cells, the surface recombination must be mitigatedby surface passivation. In general, there are two types of surface passivation schemes. One is“chemical passivation,” which is the termination of the active surface defects such as Si dangling
bonds to form Si-H and Si-O bonds by hydrogen treatment and oxidation of the Si surface,respectively, resulting in the reduction of interface states. Another type of surface passivation is“field-effect passivation, " which is the carrier population controlled by the electric fields near thesurface, originating from the charged ions or fixed charges in a surface layer. This electric field
increases the concentration of one type of carrier and decreases another type of carrier at thesurface, causing significantly different concentrations between these two types of carriers. Thissignificant difference in the two types of carrier concentration at the surface allows the reductionof surface recombination even though there are defect states at the interface. For instance, the20CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

positively charged silicon nitride (SiNy) is commonly used as a surface passivation layer due toits excellent surface passivation quality originating from both chemical and field-effectpassivation. However, the layer must be locally etched, usually performing fire-through withscreen-printed Ag electrode to contact metal electrode. Hence, the front surfaces are locallycontacted with metal electrodes in the conventional c-Si solar cell with Al-BSF, PERC, and
TOPCon structures. The SRV at the metal contacted area is extremely large (>105 cm s−1), and J0
will be unacceptably high as ~2000 fA cm−2 when using 2 Ω cm c-Si covered with only 1% of themetal electrode [12]. Thus, carrier-selective passivating contacts would realize excellent surface
passivation with a high conductivity without direct contact with metal electrodes. It would be abreakthrough technology to access a practical limit of solar cell performance. So far, SHJ solar
cells have achieved the highest conversion efficiency (η) of 26.81% [13], with carrier-selective
passivating contacts on both sides for p- and n-type contacts.

2.2.2

Theory of carrier-selective contacts

Here, the theory of carrier-selective contacts is explained by referring [14]. In a semiconductormaterial, the two types of carriers (electrons and holes) are simultaneously affected by two forces.
One is an electric force, and the other is a generalized force related to the gradient of the carrier
densities.
In the case of assuming only electric field (E) exists in the system, the charge current of particlek (Jf,k, electron is k = e or hole is k = h) is given by
𝑗f,k = 𝑞𝑛k 𝜇k 𝐸 ,(2.46)

𝑞𝑛k 𝜇k 𝐸 = 𝜎k 𝐸 ,(2.47)𝜎k 𝐸 = −𝜎kgrad(𝑧k 𝑞𝜑) .
𝑧k 𝑞(2.48)

Where q is the elementary charge, nk is the charge carrier density, μk is the mobility of the charge
carrier, φ is the electric potential, and zkq is the charge.
In the case of assuming only a concentration gradient of charge carriers k exists without anelectric field, their charge current is given by Fick’s law.

21

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS(2.49)

𝑗d,k = 𝑧k 𝑞(−𝐷k grad 𝑛k )𝑧k 𝑞(−𝐷k grad 𝑛k ) = −𝑧k 𝑞 𝑛k 𝐷kgrad 𝑛k
𝑛k(2.50)Take into consideration of Einstein’s relation of𝐷k =𝜇k 𝑘B 𝑇,𝑞

(2.51)

hence,−𝑧k 𝑞 𝑛k 𝐷kgrad 𝑛k𝑧k 𝜎k 𝑘B 𝑇 grad 𝑛k
=−
𝑛k
𝑞
𝑛k

(2.52)

is given. By applying the chemical potential of charge carriers k with a standard value of thechemical potential (Φchem,k,0), their charge current is given by

𝑗d,k = −𝜎kgrad 𝛷chem,k .
𝑧k 𝑞(2.53)In the actual case, the electrical force and chemical force are present simultaneously, and the total
current must be considered as an electrochemical potential of charge carriers (ηk). Hence, theresulting force can be described as−grad (𝛷chem,k + 𝑧k 𝑞𝜑) = −grad 𝜂k .

(2.54)The electrochemical potential means a quasi-Fermi energy of electron (EFn) and a quasi-Fermi
energy of hole (EFp), which describes the occupation of states in the conduction band or thevalence band. Hence, they can be described as𝜂e = 𝐸Fn and 𝜂h = − 𝐸Fp ,(2.55)

ze of −1 is for the charge current of electrons in all cases. Therefore, the charge current of electrons
is given by
22

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

𝑗e =

𝜎e
𝜎e
grad 𝜂e =grad 𝐸Fn .
𝑞𝑞(2.56)

ze of 1 is for the charge current of holes in all cases. Therefore, the charge current of holes is givenby𝑗h = −𝜎h𝜎hgrad 𝜂h =grad 𝐸Fp .
𝑞
𝑞(2.57)

Eqs. (2.56) and (2.57) describe the only actual charge currents of electrons and holes insemiconductors in a steady state. In the c-Si solar cells with carrier-selective passivating contact
layers for n-type and p-type, the Voc is defined as an energy difference between EFn and EFp.
Therefore, to obtain high Voc, the energy drop must be mitigated. As described in Eqs. (2.56) and(2.57), the energy drop can be mitigated by increasing conductivity and reducing one type ofcarrier concentration for each carrier type.

Fig. 2.1

Band diagram of the c-Si solar cell with carrier-selective contacts under open-circuit

conditions.

23

CHAPTER 2

2.2.3

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

Carrier selectivity

Here, the definition of carrier selectivity is explained by referring to [12]. As described above, thecarrier-selective contacts enable one polarity of charge carriers to pass to the metal electrodethrough the carrier-selective layers and to block the other polarity of carriers. Hence, the carrier
selectivity is related to the resistance to each charge carrier. The resistance for minority carriers
can be described as

𝜌m =

𝑉
.
𝑞𝑅(𝑉)

(2.58)

Where R(V) is the recombination rate, V is the internal quasi-Fermi level splitting at the contact.
The resistance for the majority carriers can be described as(2.59)

𝜌M = 𝜌c .
Where ρc is the contact resistivity. Thus, the ratio of the resistance is given by𝜌m𝑉== 𝑆(𝑉) 𝑔(𝑉) .
𝜌M 𝑞𝑅(𝑉)𝜌c(2.60)

S(V) is the selectivity of contact, which is given by
𝑉th

𝑆(𝑉) =

.

𝑞𝑅(𝑉)
𝜌c [
]
𝑉
exp (𝑉 ) − 1th(2.61)

Where Vth is the thermal voltage. g(V) is the voltage-dependent factor, which is given by𝑔(𝑉) =𝑉𝑉
𝑉th [exp (𝑉 ) − 1].

th

The selectivity factor is defined as a logarithmically scaled selectivity as follows.

24(2.62)

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

𝑆10 (𝑉) = log10 [𝑆(𝑉)](2.63)As described below, the recombination rate of the contact can be expressed as an ideal Shockley
diode equation.

𝑅(𝑉) =𝐽0
𝑉[exp ( − 1)]𝑞
𝑉th(2.64)

Where J0 is the emitter saturation current density as a recombination current density parameter.
The ideality factor is one even if Auger recombination dominates due to a highly doped surface
region in low injection. The ρc is often independent of voltage. Therefore, selectivity can be
expressed as

𝑆=

𝑉th
.
𝜌c 𝐽0

(2.65)

Hence, the selectivity of the S10 value can be described as
𝑉th
),
𝜌c 𝐽0

(2.66)

𝑉th),𝜌c,e 𝐽0,e(2.67)

𝑉th).
𝜌c,h 𝐽0,h(2.68)

𝑆10 = log10 (for the electron-selective contact,

𝑆10,e = log10 (for the hole-selective contact,

𝑆10,h = log10 (25

CHAPTER 2

2.2.4

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

Passivating contact materials and structures

Therefore, to obtain excellent carrier selectivity, conductivity for only one polarity type must be
increased for one side. The band alignments between Si and the carrier-selective materials should
be carefully considered. Hence, the band offsets and work functions of the materials with c-Si areengineered to achieve excellent carrier selectivity. The band offsets and work functions dependon the material itself, but the composition and structure of each material affect their properties.
Figure 2.2 shows the band diagrams of Si surfaces with several types of carrier-selective contacts
for the n-type c-Si under open-circuit conditions. Type 1 indicates that hole-selective contact as ap+ doping layer is introduced by a boron (B) diffusion layer on the Si surface. Type 2 indicatesthat an external potential source, such as a deposited layer with negative fixed charges, modifies
the hole-selective contacts. For instance, the highly negatively charged Al2O3 layer demonstrated
a lower sheet resistance and a possibility of usage as a p+ emitter for the n-type c-Si solar cells[15]. Type 3 indicates the hole-selective contacts formed by a heterojunction with a wide-bandgap
material, realizing a suitable Fermi level alignment with c-Si and blocking electrons by the large
conduction band edge. Type 4 indicates the hole-selective contact by the surface band bending
induced by a heavily doped p-type semiconductor such as B-doped amorphous Si (p-a-Si). This
type of structure is conventionally used for SHJ solar cells. Type 5 indicates the hole-selective
contact by increased hole concentration at the c-Si surface induced by the work function
difference between materials such as MoO3, V2O5, and WO3 [16]. Type 6 indicates electronselective contact by increased electron concentration at the c-Si surface induced by the highphosphorus (P)-doped poly c-Si and an insertion of the tunnel SiO2 layer, which has a large
valence band offset. This type of carrier-selective contact includes TOPCon.
As described above, many types of carrier-selective contact structures exist, but in this research,
the external charge-induced carrier-selective contact (type 2) was used. Because it does not
require the dopants and the diffusion process using high-temperature annealing in a specific gas
atmosphere such as POCl3 and BBr3. In addition, as described in Eq. (2.34), surface hole density
is proportional to squared surface charge. Hence, increasing the surface charge density cansignificantly increase hole selectivity. Additionally, as described in Eq. (2.35), since the Seff isinversely proportional to the surface charge density, the surface passivation quality can besignificantly improved. Importantly, as described in Eq. (2.35), the surface charge density
significantly affects Seff compared with the Nit. Hence, the field-effect passivation would be a
relatively significant solution compared with a chemical passivation to reduce Seff. The reported
carrier-selective contacts of the J0 and ρc are plotted in Figure 2.3 with S10 values [17–29]. There
has been developed a lot of carrier-selective contacts but the example of the external chargeinduced carrier-selective contact is few.
26

CHAPTER 2

Fig. 2.2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS(a) Type 1(b) Type 2(c) Type 3(d) Type 4(e) Type 5(f) Type 6

Band diagrams of the c-Si solar cell with several types of carrier-selective contacts

under open-circuit conditions.
27

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

Fig. 2.3 Carrier selectivity for the reported carrier-selective contacts.

2.3 Summary
Since surface recombination significantly affects the solar cell performance, the surface
recombination must be mitigated by surface passivation. One of the critical solutions to accessthe theoretical limit of c-Si solar cells is carrier-selective passivating contacts. The carrier
population control enables the reduction of surface recombination and the increase of conductivity
for one type of carrier. There are many types and materials of carrier-selective passivating contacts,
but the surface charge-induced type is highly expected to be the best choice because of thesignificant effect on surface recombination according to the simplified SRH formalism.

28

CHAPTER 2

SURFACE RECOMBINATION THEORY AND PASSIVATING CONTACTS

References[1]

T. Trupke, M. A. Green, P. Wurfel, P. P. Altermatt, A. Wang, J. Zhao, and R. Corkish,“Temperature dependence of the radiative recombination coefficient of intrinsic crystalline
silicon,” J. Appl. Phys. 94, 4930 (2003).

[2]

M. J. Kerr and A. Cuevas, “General parameterization of Auger recombination in crystalline
silicon,” J. Appl. Phys. 91, 2473 (2002).

[3]

A. Richter, S. W. Glunz, F. Werner, J. Schmidt, and A. Cuevas, “Improved quantitativedescription of Auger recombination in crystalline silicon,” Phys. Rev. B 86, 165202,(2012).

[4]

W. Shockley and W. T. Read Jr., “Statistics of the recombinations of holes and electrons,”
Phys. Rev. 87, 835 (1952).

[5]

R. Hall, “Electron-Hole Recombination in Germanium,” Phys. Rev. 87, 387 (1952).

[6]

A. S. Grove and D. J. Fitzgerald, “Surface effects on p-n junctions: Characteristics ofsurface space-charge regions under non-equilibrium conditions,” Solid. State. Electron. 9,783 (1966).

[7]

D. J. Fitzgerald and A. S. Grove, “Surface recombination in semiconductors,” IEEE
Trans. Electron Devices 15, 426 (1968).

[8]

K. R. McIntosh and L. E. Black, “On effective surface recombination parameters,” J. Appl.
Phys. 116, 014503 (2014).

[9]P. Saint-Cast, E. Billot, P. Olwal, S. Kühnhold, A. Richter, M. Hofmann, J. Rentsch, and R.
Preu, “Surface passivation of highly and lowly doped p-type silicon surfaces with PECVD
Al2O3 for industrially applicable solar cell concepts,” In: Proceedings of the 26th European
PV Solar Energy Conference and Exhibition, Hamburg, Germany, pp. 1125–1128 (2011).

[10]

A.G. Aberle, S.W. Glunz, A.W. Stephens, and M.A. Green, “High efficiency silicon solar
cells: Si/SiO2 interface parameters and their impact on device performance,” Prog.
Photovolt. 2, 265 (1994).

[11]

L. E. Black, “New perspectives on surface passivation: Understanding the Si–Al2O3
interface,” Ph.D. thesis, Australian National University (2015).

[12]

J. Schmidt, R. Peibst, and R. Brendel, “Surface passivation of crystalline silicon solar cells:
Present and future,” Sol. Energy Matter. Sol. Cells 187, 39 (2018).

[13]

H. Lin, M. Yang, X. Ru, G. Wang, S. Yin, F. Peng, C. Hong, M. Qu, J. Lu, L. Fang, C. Han,
P. Procel, O. Isabella, P. Gao, Z. Li, and Xi. Xu, “Silicon heterojunction solar cells with up
to 26.81% efficiency achieved by electrically optimized nanocrystalline-silicon holecontact layers,” Nat. Energy 8, 789 (2023).
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U. Wurfel, A. Cuevas, and P. Wurfel, “Charge carrier separation in solar cells,” IEEE J.
Photovolt. 5, 461 (2014).

[15]

D. Choi and K. Seo, “Field-induced radial junction for dopant-free crystalline silicon
microwire solar cells with an efficiency of over 20%,” Adv. Energy Mater. 11, 2003707(2021).

[16]

L. G. Gerling, S. Mahato, C. Voz, R. Alcubilla, and J. Puigdollers, “Characterization oftransition metal oxide/silicon heterojunctions for solar cell applications,” Appl. Sci. 5, 695,
(2015).

[17]

D. Hiller, P. Hönicke, and D. König, “Material combination of tunnel-SiO2 with a(sub-)monolayer of ALD-AlOx on silicon offering a highly passivating hole selective
contact,” Sol. Energy Mater. Sol. Cells 215, 110654 (2020).

[18]

R. Gogolin, D. Zielke, W. Lövenich, R. Sauer, and J. Schmidt, “Silicon Heterojunction
Solar Cells Combining an a-Si:H (n) Electron-collector with a PEDOT:PSS Holecollector,” Energy Procedia 92, 638 (2016).

[19]

G. Nogay, J. Stuckelberger, P. Wyss, E. Rucavado, C. Allebéb, T. Koida, M. MoralesMasisa, M. Despeisse, F.-J. Haug, P. Löper, and C. Ballif, “Interplay of annealing
temperature and doping in hole selective rear contacts based on silicon-rich silicon-carbide
thin films,” Sol. Energy Mater. Sol. Cells 173, 18 (2017).

[20]

H. Lin, M. Yang, X. Ru, G. Wang, S. Yin, F. Peng, C. Hong, M. Qu, J. Lu, L. Fang, C. Han,
P. Procel, O. Isabella, P. Gao, Z. Li, and Xi. Xu, “Silicon heterojunction solar cells with up
to 26.81% efficiency achieved by electrically optimized nanocrystalline-silicon holecontact layers,” Nat. Energy 8, 789 (2023).

[21]

J. Bao, C. Chen, L. Ma, C. Huang, X. Zhang, Z. Wang, C. Chen, S. Zhan, R. Liu, Z. Qiao,
Z. Du, Z. Liu, and J. Chen, “Towards 24% efficiency for industrial n-type bifacial
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European Photovoltaic Solar Energy Conference and Exhibition, online, pp. 160–164(2020).

[22]

D. Yan, A. Cuevas, Y. Wan, and J. Bullock, “Passivating contacts for silicon solar cells
based on boron-diffused recrystallized amorphous silicon and thin dielectric interlayers,”
Sol. Energy Mater. Sol. Cells 152, 73 (2016).

[23]

Y. Wan, C. Samundsett, J. Bullock, T. Allen, M. Hettick, D. Yan, P. Zheng, X. Zhang, J.
Cui, J. McKeon, A. Javey, and A. Cuevas, “Magnesium fluoride electron-selective contacts
for crystalline silicon solar cells,” ACS Appl. Mater. Interfaces 8, 14671 (2016).

[24]

D. L. Young, W. Nemeth, S. Grover, A. Norman, B. G. Lee, and P. Stradins, “Carrierselective, passivated contacts for high efficiency silicon solar cells based on transparent
conducting oxides,” Energy Procedia 55, 733 (2014).
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X. Yang, P. Zheng, Q. Bi, and K. Weber, “Silicon heterojunction solar cells with electron
selective TiOx contact,” Sol. Energy Mater. Sol. Cells 150, 32 (2016).

[26]

S. W. Glunz, F. Feldmann, A. Richter, M. Bivour, C. Reichel, H. Steinkemper, J. Benick,and M. Hermle, “The irresistible charm of a simple current flow pattern–25% with a solar
cell featuring a full-area back contact,” In: Proceedings of the 31st European Photovoltaic
Solar Energy Conference and Exhibition, Hamburg, Germany, pp. 259–263 (2015).

[27]

S. Y. Herasimenka, W. J. Dauksher, and S. G. Bowden, “>750 mV open circuit voltage
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[28]

M. Rienäcker, M. Bossmeyer, A. Merkle, U. Römer, F. Haase, J. Krügener, R. Brendel, and
R. Peibst, “Junction resistivity of carrier-selective polysilicon on oxide junctions and its
impact on solar cell performance,” IEEE J. Photovolt. 7, 11 (2017).

[29]

R. Brendel and R. Peibst, “Contact selectivity and efficiency in crystalline silicon
photovoltaics,” IEEE J. Photovolt. 6, 1413-1420 (2016).

31

CHAPTER 3

CHARACTERIZATION TECHNIQUES

32

CHAPTER 3

CHARACTERIZATION TECHNIQUESCHAPTER 3
CHARACTERIZATION TECHNIQUES

3.1 Carrier lifetime measurements
The carrier lifetime can be determined by measuring the photoconductance of the semiconductor.
The photoconductance can be analyzed by measuring the eddy current generated in an inductively
coupled coil or the reflectance of microwaves. Hence, the quasi-steady state photoconductance(QSSPC) method and the microwave photoconductivity decay (µ-PCD) method have been widelyused for Si-based devices.

3.1.1

Quasi-steady state photoconductance (QSSPC)

The quasi-steady state photoconductance (QSSPC) method was first proposed by Sinton andCuevas [1], which has been widely used to evaluate the carrier lifetime of c-Si solar cells. This
technique can obtain the Δn by analyzing the conductivity of the light-illuminated sample throughan inductively coupled coil. This method allows us to obtain injection-dependent carrier lifetimes
because the excitation intensity gradually decays over time using a Xenon flash lamp. The shortwavelength light must be cut using a long pass filter as a homogeneous generation must occur inthe entire wafer. The gradual decay of the Δn over time can be described as33

CHAPTER 3

CHARACTERIZATION TECHNIQUES

𝑑𝛥𝑛(𝑡)= 𝐺(𝑡) − 𝑈(𝑡) .
𝑑𝑡

(3.1)

Where G(t) is the generation rate at the time of t, and U(t) is the recombination rate at the time oft. By substituting the relationship between the recombination rate and the effective lifetime
described as

𝑈=

𝛥𝑛
.
𝜏eff

(3.2)

The effective lifetime as a generalized analysis [2] can be determined by𝜏eff =𝛥𝑛.
𝑑𝛥𝑛(𝑡)
𝐺(𝑡) −𝑑𝑡(3.3)

The generation rate (G) can be analyzed by the photocurrent of a reference solar cell, which ismounted in the equipment near the sample on the sample stage. In addition, the sample's opticalproperty must be considered an optical factor (OF). It is a ratio of optical properties between asample and a reference cell. In this study, the OFs were determined by the calculated photon
current densities from the absorbance spectra. For the flat Si substrates, the OF should be 0.7, but
the wafer with surface roughness or ARC must apply suitable OF values for each case. The
generation rate at the time of t can be described as

𝐺(𝑡) = 𝑂𝐹

𝐼ref (𝑡)
.
𝑞𝑊ref

(3.4)

Where Iref(t) is the photocurrent of a reference solar cell at the time of t, and Wref is the wafer
thickness of the reference cell. The injection level can be determined from the equation below.

𝛥𝑛(𝑡) =𝛥𝜎(𝑡)𝑞(𝜇n + 𝜇p )𝑊(3.5)

Where Δσ is the measured excess conductivity in the sample, μn and μp are the electron and holemobilities, and W is the wafer thickness.
At last, the effective lifetime can be calculated using Eq. (3.3). In addition, implied open-circuit
34

CHAPTER 3

CHARACTERIZATION TECHNIQUES

voltage (iVoc) can be determined at each illumination intensity using the equation below.

𝑖𝑉oc =(𝑁dop + 𝛥𝑛)𝛥𝑛
𝑘B 𝑇
ln [
]
𝑞
𝑛i 2

(3.6)

Where kB is the Boltzmann constant, T is the absolute temperature, Ndop is the doping
concentration, and ni is the intrinsic carrier density. Moreover, the J0 can be derived from the
below equation of the Kane and Swanson method [3] using the auger-corrected effective minority
carrier lifetime (τcorr) over the high injection level.
2𝐽0 (𝑁dop + 𝛥𝑛)
1
1
=
+
.
𝜏eff 𝜏bulk
𝑞𝑛i 2 𝑊

(3.7)

At high injection levels, since the bulk lifetime dominates the Auger lifetime, the equation can be
described as
2𝐽0 (𝑁dop + 𝛥𝑛)
1
1
=
+
.
𝜏eff 𝜏Auger
𝑞𝑛i 2 𝑊

(3.8)

Hence, J0 can be derived from the slope in the relation between τcorr and injection level. The
schematic illustration of the lifetime tester equipment is illustrated in Figure 3.1. In this work,
Sinton lifetime tester WCT-120 was used for τeff and iVoc analysis.

Fig. 3.1 Schematic illustration of setup for the QSSPC measurements.

35

CHAPTER 3

3.1.2

CHARACTERIZATION TECHNIQUESMicrowave photoconductivity decay (µ-PCD)

The microwave photoconductivity decay (µ-PCD) utilizes pulsed laser (~200 ns) light for theexcitation of carriers in the substrate. The photoconductivity is indirectly analyzed by detecting
microwave reflection. The excited carriers increase the conductivity of the material, and thephotoconductivity increases microwave reflection. Hence, by detecting the decay of the
microwave reflection over time, an effective minority carrier lifetime (τeff) can be derived fromthe decay of excess carriers [4]. In general, the carrier lifetime can be defined as the time constantat 1/e of the maximum microwave intensity just after being irradiated by the pulsed laser light. In
this experiment, a wavelength of 904 nm was used for the excitation by pulsed laser light. Since
the system is applied pulsed laser excitation, the µ-PCD method cannot measure Δn dependence
on the τeff. Since the usual areal photon density is 5×1013 cm−2, assuming a homogeneous
generation in the entire substrate with a thickness of 280 µm, the instant Δn would be 1.2×1015
cm−3. Hence, this method has a transient feature: the generation rate (G) equals zero. Therefore,
Eq. (3.3) can be described as𝜏eff =𝛥𝑛.
𝑑𝛥𝑛(𝑡)−𝑑𝑡(3.9)

The deferential Δn has a solution of Δn = A exp(−t/τeff). The τeff can be derived from the fitting of
the single exponential curves. Additionally, the unique important feature of this method ismeasurable τeff mappings. The excitation area is ~1 mm2, and the movable sample stage allows us
to obtain τeff mappings. The τeff mapping can display the distribution within the sample area,
allowing us to evaluate the evenness of surface passivation quality of semiconductor materials.
This research used Kobelco Research Institute, Inc., LTA-1510EP to acquire τeff mappings by µPCD method. In this equipment, the wavelength of 349 nm is also mounted with variable laser
light intensity in the range of 5×1012 –1.1×1015 cm−2. This research used this option to acquire thepulsed laser light intensity dependence on the τeff. However, since the Si material has a large
absorption coefficient of light with a wavelength of 349 nm (~1×106 cm−1), such a shortwavelength of light is absorbed in the surface of Si substrates. Thus, homogeneous generation
cannot be realized, but the bulk quality of the material is exceptionally high enough compared tothe surface. This method can also evaluate the surface passivation quality. However, the carrier
generation almost occurred in the space charge region. The minority carrier density is relatively
higher than the long wavelength laser excitation, which might result in lower τeff. Moreover,steady-state light can be illuminated during the measurement as a bias light in this system. In this36

CHAPTER 3

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research, the transient lifetime was also evaluated under the bias light illumination, varying thebias light intensity in the range of 0.025–1.2 suns. In general, if the band bending exists at thesurface of the material, the band bending is reduced by the light illumination due to the increaseof Δn in the space charge region. If the surface passivation is dominated by field-effect passivation,
the surface passivation quality would be gradually decreased with increasing the steady-state light.
Hence, this research utilizes this method to elucidate the surface passivation scheme.

Fig. 3.2 Schematic illustration of setup for the µ-PCD measurements.

3.2

X-ray photoelectron spectroscopy (XPS)

Principle
X-ray photoelectron spectroscopy provides information about the chemical bond states, theatomic composition at the surface of materials, and the thickness of the ultrathin layer on thesubstrates. When the X-ray hits the surface of materials, the generated photoelectrons are emitted
from the surface with kinetic energies corresponding to each material and chemical bond states.
The relation between the kinetic energy and the binding energy of the chemical bond states is
described as
Ek = hν– Eb −Φ.

(3.10)

Where Ek is the kinetic energy of photoelectrons, hν is the X-ray energy, Eb is the binding energyof the chemical bond states, and Φ is the spectrometer work function. Since the hν and Φ are
known values, if Ek is measured, Eb can be determined. In the XPS measurement, Ek and theintensity are acquired simultaneously. Concretely, the intensity is defined as the detected numberof photoelectrons with each Ek. Hence, XPS measurement acquires the spectrum as a function of
the binding energy. The intensity peaks are shifted to higher or lower binding energy. The
37

CHAPTER 3

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chemical shift allows us to know the information about the chemical bond states. For instance, inthe case of the SiO2/Si interface, the SiO2 thin layer does not entirely consist of “SiO2”. Usually,
SiO2 consists of four chemical bond states such as Si1+, Si2+, Si3+, and Si4+(SiO2). With increasing
the amount of bonded oxygen, the silicon oxide peak can be observed at the higher binding energy.
According to the literature [5], the relative binding energies to the Si 0 (the bulk of Si) for eachchemical bond state are 0.95, 1.75, 2.48, and 3.90 eV, respectively. The components of Si1+, Si2+,
and Si3+ are known as sub-oxides, which are lower-quality SiO2. Therefore, the film quality canbe evaluated by analyzing the sub-oxides and SiO2 components. Thus, high-resolution analysis isessential to evaluate such chemical shifts. This work used Al Kα (hν = 1486.6 eV) as an X-ray
source to generate photoelectrons. For the detection of photoelectrons, they were collected froma normal angle; in other words, detected photoelectrons perpendicularly emitted from the surface
by setting a take-off angle of 90°. The reason for detectable photoelectrons from less than ~10 nm
depth is due to the decay of photoelectrons by the inelastic scattering in the surface of materials.
Hence, the XPS measurement allows us to do very surface-sensitive analysis and has a massive
advantage for the ultrathin layer, like an Al-doped SiOx layer formed in this study. The
hemispherical analyzer enables us to obtain high-resolution binding energies. However, during
the measurement on the insulator surface, the charge-up occurs at the X-ray irradiated surface,leading to an unexpected shift to the higher binding energy. Therefore, the peaks must be
calibrated to evaluate the chemical shifts accurately. In general, all acquired peaks are calibrated
by the acquired C 1s peak positions. However, in this research, the chemical shift concerning theSi0 is not crucial. Hence, the peaks of Si0 were fixed to the reference value of 99.2 eV [6]. The
schematics of the XPS system (Shimadzu Corporation, Kratos AXIS-ULTRA DLD) used in thisstudy are shown in Figure 3.3.

Fig. 3.3 Schematic illustration of XPS measurements.

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CHAPTER 3

CHARACTERIZATION TECHNIQUES

Evaluation of the surface atomic concentration
The surface atomic composition can also be obtained by XPS analysis. In this study, I acquired
all core-level spectra of each atom, Si 2p, O 1s, and Al 2p. To evaluate the atomic concentrationof the layer, the areas of peak fitted intensities for each atom are applied to the equation below
[7],
𝐼i
(𝑅𝑆𝐹
)
i
𝑋i =
.
𝐼j
n
∑j=1 ()
𝑅𝑆𝐹j(3.11)

Where Ii is the area of intensity for the specific atom, RSFi is the relative sensitivity factor for the
specific atom, Ij is the area of intensity for each atom, and RSFj is the relative sensitivity factor
for each atom. The RSF of 0.283 for Si 2p [8], 0.711 for O 1s [8], and 0.193 for Al 2p [8] wereapplied to evaluate the surface atomic concentration of oxidized Al and oxidized Si. Each area ofintensity can be derived from the peak-fitted spectra by subtracting the background using the
Shirley–Proctor–Sherwood method [9, 10] and peak fitting using the Voigt function.
Evaluation of the SiOx thickness
The SiOx thickness can be evaluated from the area of intensity for Si0 (ISi) and the sum of the areaof intensity for the components of 𝐼SiO2 using the equation below [5],𝑑 = 𝑙SiO2 ln (𝐼SiO2𝑛Si 𝜎Si 𝑙Si).
𝐼Si 𝑛SiO2 𝜎SiO2 𝑙SiO2(3.12)

Where 𝜎Si and 𝜎SiO2 are the photoionization cross sections, 𝑛Si and 𝑛SiO2 are the density of
Si atoms, 𝑙Si and 𝑙SiO2 are the escape depths for Si and SiO2, respectively. 𝑛Si of 5.00×1022
cm−3, 𝑛SiO2 of 2.28×1022 cm−3, 𝑙Si of 2.7 nm [11], 𝑙SiO2 of 3.7 nm [11], and the ratio of thephotoionization (𝜎SiO2 /𝜎Si) of 1.1 [12] were used in this research.

39

CHAPTER 3

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3.3 Current–voltage (I–V) measurement
Current and voltage measurement (I–V) is the method to characterize the relationship of current
and voltage, which are applied between the top and bottom metal electrodes on the device
structure, such as solar cells, and the more fundamental structure of metal oxide semiconductor(MOS). While sweeping the bias voltage, the current is measured. In this study, I used a precision
semiconductor parameter analyzer (Hewlett-Packard Company, 4156 A) mounted with a kelvintype four-terminal connection system and acquired dark I–V characteristics for MOS structure.
The current densities were calculated by the measured current values divided by the top electrodearea.
Tunneling current density
Then, the calculated dark current density on the MOS structure is defined as a tunneling current
density in this work. When the negative bias voltage is applied on top of the electrode, the n-type
Si band is bent upward, and an inversion layer is created. Hence, the hole current is dominant as
a tunneling current through the oxide layer at the inversion condition. On the other hand, when apositive bias voltage is applied on top of the electrode, the n-type Si band is bent downward, andthe electrons are accumulated at the Si surface. Thus, the electron current dominates as a tunnelingcurrent through the oxide layer at the accumulation condition. As considered in this theorem, thetunneling current density should be evaluated for the hole tunneling current density under
applying the negative bias voltage. Hence, in this study, I acquired a current density at theinversion condition to evaluate the hole tunneling current density. The depiction of the MOS
structure and the band structures at each voltage-applied condition is depicted in Figure 3.4. In
this configuration, a similar work function material should be selected as a top electrode to theelectron affinity of n-type Si to avoid creating a Schottky diode by the work function difference.
Thus, in this work, Al was used for the top electrode with a work function of 4.28 eV [13].

40CHAPTER 3

CHARACTERIZATION TECHNIQUES

Fig. 3.4 Depiction of the MOS and band structures when positive and negative bias are applied.

41

CHAPTER 3

CHARACTERIZATION TECHNIQUES

Characteristics of solar cell performance
I–V measurements can also evaluate the solar cell performance. The device structure must be
fabricated with a finger or grid electrode and the bottom electrode on a Si substrate with a p-njunction. The test is performed by measuring the I–V characteristics under 1 sun illumination. The
measurement must undergo a short-circuit condition (at zero bias), open-circuit condition (at zerocurrent), and between these bias conditions to determine the short-circuit current density (Jsc), Voc,
and maximum power point (Pmax) to derive fill factor (FF), and η. The typical photo J–V curve isshown in Figure 3.5 with Jsc, Voc, and maximum power point (Pmax).

Fig. 3.5 Typical photo J–V curve with Jsc, Voc, and Pmax.

Also, the formula to evaluate the FF and the η are described as
𝑃max,
𝐽sc 𝑉oc(3.13)𝑃max
𝐽sc 𝑉oc FF=.
𝑃input𝑃input(3.14)

𝐹𝐹 =

𝜂=

42

CHAPTER 3

CHARACTERIZATION TECHNIQUES

Contact resistivity (Cox and Strack analysis) [14]The total resistance (Rtotal) can be described as the sum of the spreading resistance (Rspread), thecontact resistance (Rc), and the residual resistance due to a substrate or the contact resistance ofthe back side contact (R0).
𝑅total = 𝑅spread + 𝑅c + 𝑅0

(3.15)

By using ρc, Rc can be described as

𝑅c =

𝜌c
1 2
𝜋𝑑
4

=

𝜌c
.
𝐴

(3.16)

Where d is the diameter of the circular electrode, and A is the area of the electrode. Hence, the ρccan be derived from the slope of the (Rtotal − Rspread) as a function of the inverse area of the electrode.

𝑅total − 𝑅spread =

𝜌c
+ 𝑅0
𝐴

(3.17)The Rspread is determined from the below equation.
𝜌𝐵
2𝑑(3.18)

24𝑡
arctan ( )𝜋𝑑(3.19)

𝑅spread =𝐵≈

Where B is the geometrical factor, which corrects for the finite thickness of the substrate, ρ is thesubstrate resistivity, and t is the substrate thickness. This analysis is known as the Cox and Strack
method, and the Cox and Strack equation is given by

𝑅total =𝜌
4𝑡𝜌c
arctan ( ) + + 𝑅0 .
𝜋𝑑𝑑𝐴

(3.20)

Hence, using this equation, the ρc can be analyzed from the measured I–V data for several sizesof electrodes. This study used 13 circular electrodes ranging from 0.28 to 32.15 mm2.
43

CHAPTER 3

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3.4 Capacitance–voltage (C–V) measurement
The capacitance and voltage (C–V) measurement [15] acquires capacitance while sweeping thevoltage applied to the MOS structure. This technique widely analyzes interface properties such as
Dit, Qf, dox, and εox. The applied voltage contains an AC component, and 1 MHz was applied forthis study to avoid responding to the minority carriers. To obtain the Qf value, the flat band voltage(VFB) was estimated from the flat band capacitance (CFB). By using Debye length (LD), the CFB is
described as
1
1
𝐿D
=
+
.
𝐶FB 𝐶ox 𝜀Si(3.21)

Where Cox is the oxide capacitance, the Cox can be measured at an accumulation condition. The
LD is given by𝜀Si 𝑘B 𝑇
𝐿D = √ 2.
𝑞 𝑁A(3.22)

Then, the VFB was derived from the measured C–V curves at the voltage corresponding to thecalculated CFB. The Qf value can be calculated using the equation below.

𝑄f =

𝐶ox
(𝛥𝛷ms − 𝑉FB )𝑞𝐴(3.23)𝐴 𝜀r 𝜀0𝑑(3.24)

𝐶ox =Where ΔΦms is the work function difference between metal and Si, A is the electrode area, and εris the relative permittivity of the insulator.
Mott-Schottky plot [16]The built-in potential (Vbi) of the semiconductor can be analyzed by Mott-Schottky plot analysis.
The Mott-Schottky plot is the plot of inverse squared capacitance (1/C2) concerning the applied
voltages. The plots are extrapolated from 0 V to the interception of the horizontal axis.
44

CHAPTER 3

CHARACTERIZATION TECHNIQUESConsidering the relation between the 1/C2 and the voltage described below, the Vbi can be derived
from the interception of the voltage axis,
1
2
𝑘B 𝑇
=
(𝑉 − 𝑉bi −).
22𝐶
𝜀𝜀0 𝐴 𝑞𝑁dop
𝑞(3.25)

Where ε is the relative permittivity of the material, ε0 is the permittivity of vacuum, respectively.
When 1/C2 = 0, the Vbi can be derived as𝑉bi = 𝑉 −

𝑘B 𝑇
.
𝑞

(3.26)

Hence, Vbi can be obtained using the Mott-Schottky plot and some known parameters: kB, T, andq.

3.5 Four-point probe techniqueFour-point probe technique determines the sheet resistance (Rsheet) of the material. As shown inthe schematics of the four-point probe technique in Figure 3.6, during the measurement, while asmall current passes through the two outer probes, voltage is measured between the other two
inner probes. Hence, the sheet resistance can be obtained from the measured voltage and the flow
current. In addition, resistivity (ρ) can be described as𝜌 = 𝑅sheet 𝑑 ,(3.27)

where d (cm) is the film thickness. This method can avoid the lead and contact resistance fromthe measurement because the probes are separated for measuring voltage and flowing current,
resulting in a more precise measurement of the Rsheet. This research used a four-point probe
technique (Mitsubishi Chemical Analytech Co., Ltd., Loresta-AX MCP-T370) in the dark and 1
sun illumination conditions.

45

CHAPTER 3

CHARACTERIZATION TECHNIQUES

Fig. 3.6 Schematics of the principle for the four-point probe technique.

In this study, to investigate the illumination effect on the Rsheet, the measurements were also
performed under 1 sun illumination. The experimental setup is illustrated in Figure 3.7. The
samples are illuminated from the front side, and the probes are placed on the rear side of thesamples to avoid the shading of probes. Considering the bulk lifetime is >10 ms, it is >10 times
longer carrier diffusion length of ~0.36 cm compared with the substrate thickness of 280 µm.
Hence, generated holes at the front surface can also reach the rear surface before therecombination. The diffusion length (L) can be described as √𝐷𝜏. Where D is the diffusion
coefficient of carriers and τ is the carrier lifetime in the bulk. In this case, D is the diffusion
coefficient for holes of 12.9 cm2 s−1 [15]. Therefore, this method is available to obtain Rsheet under
1 sun illumination.

Fig. 3.7 Schematic illustration of setup for the sheet resistance measurement under 1 sunillumination.
46

CHAPTER 3

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3.6 Fourier transform infrared spectroscopy (FTIR)Infrared (IR) spectroscopy gives us information about the interaction between IR light andmaterials. In general, IR absorption spectroscopy is applied since each chemical bond absorbs
light with specific frequencies, which is involved in their nature and the chemical environment
surrounding them. Therefore, IR spectroscopy allows us to investigate the chemical compositionof materials. Using an interferometer, it is simultaneously measurable for all the wavelengths in
an IR spectrum. Hence, the system does not require a monochromator. In this study, FTIR
measurements were performed to estimate the atomic density of the interfacial SiOx layer. The
atomic density of the SiOx can be determined by observing the wavenumber of longitudinal
optical (LO) and transversal optical (TO) phonons modes [17]. The frequencies of the TO and LO
modes are described as1√2[𝛼sin2 (𝜃/2) + 𝛽cos2 (𝜃/2)]/𝑚 ,2𝜋(3.28)

1√2{𝛼sin2 (𝜃/2) + 𝛽cos2 (𝜃/2) + 𝑍 2 𝜌/[𝜀∞ (2𝑚 + 𝑀)]}/𝑚 .
2𝜋(3.29)

𝜈TO =

𝜈LO =

Where νLO and νTO are the frequency of LO and TO modes, α and β are the central and noncentral
force constants, 𝜃 is the Si–O–Si bridging bond angle, m and M are the atomic masses of O andSi atoms, respectively, ε∞ is the permittivity at the infinite frequency of the oxide film, Z is theelectrical charge related to the movement of O atoms. Hence, SiOx atomic densities (ρ) were
calculated from the equations below,

𝜈LO 2 − 𝜈TO 2 =2𝑍 2
𝜌 = 𝐶𝜌 .
[𝜀∞ (2𝑚 + 𝑀)]𝑚

(3.30)For the estimation of C, the reported value of 1.71×10−17 cm from bulk SiO2 [18, 19] was used inthis study. However, the objective of this study is an ultrathin SiOx layer with a thickness of ~1
nm, which causes difficulty in observing off-plane vibration by the standard setup of verticalincident light. Thus, to detect the longitudinal optical (LO) phonon vibrational mode of SiOx, themeasurements were performed at an oblique incident light setting of 74° by using a variable
angular holder, which can be detected LO mode by p-polarized light due to the Berreman effect
[20].
47

CHAPTER 3

3.7

CHARACTERIZATION TECHNIQUESUV-vis-NIR spectrometryUV-vis-NIR spectroscopy analyzes the material's optical properties, such as reflectance,transmittance, and absorbance. These optical properties are acquired with varying the wavelengthof the incident light. Since the light can scatter at the surface of samples, the integrating sphere isequipped with a highly reflective coating to detect all the scattering light. The relation among thereflectance (R), transmittance (T), and absorbance (A) is described as
𝐴 = 1 − (𝑅 + 𝑇) ,(3.31)and the absorption coefficient (α) can be determined by𝐼𝑇== 𝑒 −𝛼𝑡 .
𝐼0 1 − 𝑅(3.32)

Where t is the thickness of samples.

3.8 Scanning electron microscopy (SEM)
Scanning electron microscopy (SEM) can observe surface and cross-sectional structures ofmaterials on a micrometer scale to a nanometer scale. In the microscope system, the emitted
electrons are accelerated by applying high anode bias (0.5–30 kV). Then, the accelerated electrons
are condensed into a narrow beam using the magnetic lens, which is focused on the samples. The
impinged electron beam at the sample surface can scatter back, emit secondary electrons (SE), or
generate X-rays. Generally, the surface structure is observed by collecting SE due to thegeneration of SE only at the proximity of the surface, which originated from losing their energy
at the surface by collisions. The image can be obtained by scanning the beam and collecting
electrons at each position. The above-all processes occur in a vacuum chamber with a pressure of<10−3 Pa. In this study, SEM (Hitachi, TM3030Plus) was used to observe the surface structure ofa pyramidal textured Si surface.

48

CHAPTER 3

3.9

CHARACTERIZATION TECHNIQUESTransmission electron microscopy (TEM)

Transmission electron microscopy (TEM) observes the surface and cross-sectional structures ofmaterials on a nanometer scale. The accelerated and focused electron beam is transmitted throughthe sample and detected. By detecting the intensity of transmitted electrons, it observes a spatial
distribution of the transmittance of electrons. The structures and composition of the materials
differ in electron transmittance, providing an image by acquiring transmitted electrons at eachposition. Since the electron must be transmitted, the sample must be thin enough to be less than
100 nm. The acceleration voltage is generally 100 kV, and the TEM measurement is performed
under high vacuum conditions of <10−4 Pa.

3.10

Energy dispersive X-ray spectroscopy (EDS)

Energy dispersive X-ray spectroscopy (EDS) analyses the elements' composition of the samples.
The energy of the generated X-ray is inherent to each atom. Hence, detecting the energy and theintensity gives us information about the element species and their compositions. By impinging
accelerated electrons to the samples, secondary electrons are emitted from the core level of thesample. Then, outer shell electrons are transferred to the vacancy of the inner shell to be stable,
and the excess energy corresponding to the transition energy is emitted as an X-ray. In this study,
X-rays of Si Kα, O Kα, and Al Kα were detected. The emitted secondary electron comes out fromthe K-shell, and the electron transfer occurs from L to K-shells; then, the emitted X-ray is referred
to as Kα. In general, the EDS system is mounted in SEM and TEM systems. Hence, the X-ray
intensity mapping for each atom can be obtained, referred to as EDS mapping.

3.11 Scanning probe microscopy (SPM)
Atomic force microscopy (AFM)
Scanning probe microscopy observes the surface structure of samples by scanning a physicalprobe. One of the surface imaging modes of atomic force microscopy (AFM) can produce surface
images at the atomic level. While scanning a cantilever over a sample surface, the laser light isirradiated on the cantilever, and the reflected light is detected to obtain information about thesurface structure. While probing and contacting the surface with a sharp tip on the edge of thecantilever, bending of the cantilever causes the change in the amount of the reflected laser lightdetected by a photodiode, which provides a spatial image with information of surface structure49

CHAPTER 3

CHARACTERIZATION TECHNIQUES

due to tracing of the height of cantilever. The piezoelectric element is mounted for precise
scanning and oscillates cantilever at its eigenfrequency.
Kelvin probe force microscopy (KPFM)
Kelvin probe force microscopy (KPFM) is the contact potential difference (CPD) measurement
between the tip of the cantilever and the sample surface. The CPD corresponds to a work function
difference between the two materials. Since this is the scanning probe method, it provides asurface potential mapping on the sample. When the two materials are contacted, Fermi levels areplaced at the same level. Hence, it generates contact potential (Vcpd). However, when the contact
potential is nullified by applying an external voltage, the Fermi level of the sample returns to its
original position. The applied external bias is known as backing potential (Vb). This means the Vb
equals Vcpd, as shown in Figure 3.8. Hence, the relation between CPD and work function for tip
and sample can be described as𝑉b = 𝑉CPD =Φtip − Φsample.
𝑞(3.33)

Hence, the work function of the sample can be estimated by measured Vcpd and the known work
function of the tip.

Fig. 3.8 Schematics of the principle for the KPFM measurement.

50

CHAPTER 3

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R. A. Sinton, A. Cuevas, and M. Stuckings, “Quasi-steady-state photoconductance, a
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H. Nagel, C. Berge, and A. G. Aberle, “Generalized analysis of quasi-steady-state andquasi-transient measurements of carrier lifetimes in semiconductors,” J. Appl. Phys. 86,
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[6]

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[9]D. A. Shirley, “High-resolution x-ray photoemission spectrum of the valence bands ofgold,” Phys. Rev. B 5, 4709 (1972).

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M. F. Hochella Jr. and A. H. Carim, “A reassessment of electron escape depths in silicon
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CHARACTERIZATION TECHNIQUESJ. Hölzl and F. K. Schulte, “Work Functions of Metals” in Solid Surface Physics, SpringerVerlag, Berlin (1979).

[14]

R. H. Cox and H. Strack, “Ohmic contacts for GaAs devices,” Solid-State Electron. 10,1213 (1967).

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K. Gelderman, L. Lee, and S. W Donne, “Flat-band potential of a semiconductor: Using
the Mott–Schottky equation,” J. Chem. Educ. 84, 685 (2007).

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A. Lehmann, L. Schumann, and K. Hübner, “Optical phonons in amorphous silicon oxides
Ⅰ . Calculation of the density of states and interpretation of LO － TO splittings ofamorphous SiO2,” Phys. Status Solidi B 117, 689 (1983).

[18]

Asuha, T. Kobayashi, M. Takahashi, H. Iwasa, and H. Kobayashi, “Spectroscopic andelectrical properties of ultrathin SiO2 layers formed with nitric acid,” Surf. Sci. 547, 275(2003).

[19]

C. Martinet and R. A. B. Devine, “Analysis of the vibrational mode spectra of amorphous
SiO2 films,” J. Appl. Phys. 77, 4343 (1995).

[20]

D. W. Berreman, “Infrared absorption at longitudinal optic frequency in cubic crystal
films,” Phys. Rev. 130, 2193 (1963).

52

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDECHAPTER 4
HOLE-SELECTIVE CONTACT BY ANULTRATHIN Al-DOPED SILICON OXIDETo accomplish high η of Si solar cells, the device requires interlayers with not only surface
passivation but also sufficient carrier tunnel probability. In particular, SiO2 layers have a highadvantage for the front-side passivation layer because of their extremely high optical properties
with low parasitic absorption attributed to their high bandgap properties. However, the carrier
tunneling through SiO2 decreases exponentially with increasing layer thickness [1]. In contrast, toachieve high surface passivation quality, the SiO2 layer has to be thickened by a high-temperatureannealing process, such as thermal oxidation, followed by forming gas annealing (FGA) toterminate active defects by hydrogens [2,3]. Therefore, the device performance will be limited
due to the trade-off between the carrier tunneling probability and the surface passivation quality.
Thus, this chapter overcomes this problem by impurity-doped SiOx. The film properties ofthickness and composition were investigated. In addition, surface passivation quality and holeselectivity. Moreover, Qf and interface state density were estimated by an analyzed surface
potential and simplified extended SRH formalism. The hole-selective passivating contact
mechanism is proposed by considering the important finding of the light-induced activation effect.

53

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HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

4.1 Impurity doped SiOx for carrier-selective contacts
Impurity doping into the SiO2 layer will be one of the expected solutions. Recently, D. König etal. reported that ultrathin Al–O/SiO2 stacks show high hole selectivity induced by negative fixedcharges originating from Al-induced acceptor states occupied with electrons, which are formed in
the SiO2 layer [4]. According to their proposed model, a hole-selective passivating contact —selective hole tunneling through the SiO2 layer with an excellent surface passivation — can berealized by hole tunneling via the Al-induced acceptor states and effective carrier separation due
to upward band bending of Si surface induced by the negative fixed charges existing in the SiO2
layer.
In addition, if the SiO2 layer has sufficient negative fixed charges, p-type inversion layers areformed on the n-type Si substrates. Such modulation doping emerged as a technology of providing
free carriers at the surface of a semiconductor by doping materials adjacent to the semiconductor
without direct doping of impurities in Si, which can eliminate dopant-associated defect states [5,6].
The modulation doping technology was first implemented for GaAs/AlAs in the 1970s [7]. In
recent years, Si modulation doping has emerged by using AlxGa1−xN as a dielectric layer [8], andthe SiO2 layer was used for the modulation doping of III-V semiconductors [4]. The density
functional theory (DFT) calculation revealed that Al [4], Sc, and Ga could be candidates for
impurities for modulation doping in SiO2 [9]. These impurities in SiO2 form the modulation
acceptor states and the energies of acceptor levels relative to vacuum levels are 5.36, 5.51, and5.39, respectively. Since the modulation acceptor states exist in the band gap of SiO2 and below
the Si valence-band edge, the relaxations of electrons from Si readily occur by transition into themodulation acceptor states without thermal activation, delivering a large number of holes to theSi surface at very low temperatures, which cannot be realized by the conventional doping process.

4.2 Concept of aluminum nitrate oxidation treatment
Although the ultrathin Al–O/SiO2 stacks showed an excellent hole-selective passivating contact
property, the ultrathin Al–O/SiO2 stacks were formed by rapid thermal oxidation (RTO) or
plasma-enhanced chemical vapor deposition (PECVD) for the SiO2 formation, followed by theformation of monolayer or multilayer Al–O by using atomic layer deposition (ALD). In addition,
high-temperature annealing at 1000 °C is required to activate the surface passivation [4]. These
processes lead to degraded throughput of solar cell manufacturing. On the other hand, ultrathin
SiOx layers can be simply formed by acid aqueous solutions such as HClO4 [10], HNO3 [11, 12],54

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEHCl + H2O2 [12], H2SO4 + H2O2 [12]. In addition, ozonated water [13], and direct current
anodization process can also form ultrathin SiOx layers [2]. Though these processes can be
performed at low temperatures and are thus cost-effective cell production, these SiOx layers
themselves have inferior surface passivation quality and after all require further reformation suchas thickening to ~5 nm by thermal oxidation, high-temperature annealing at 1100 °C, and FGA to
increase surface passivation quality [2, 3]. Indeed, the lowest SRV of 42 cm s−1 at a Δn of 1015
cm−3 has been reported, in which SiOx was formed using HNO3 oxidation followed by hightemperature thermal treatment [2]. However, ultrathin SiOx layers themselves with both goodsurface passivation quality and carrier-selective contact have not yet been reported by realizing
only using a low-temperature process. Therefore, this research investigated a simple wet chemical
method realizing a hole-selective passivating contact by ultrathin SiOx layers with impurity
doping.
If the wet chemical solution contains Al and the processing Si surface is terminated by the OH
group, the Al atoms are favorably adsorbed on the Si surface as a formation of Si-O-Al(OH)3(final state) due to highly reactive combinations between Si-OH and Al(OH)4 (initial state),decreasing potential energy and stabilizing [14]. The enthalpy change (ΔH) from the initial state
to the final state of −0.91 eV suggests that the final state is stabilized [14]. Al is generally
considered a metal contaminant, creating a shallow energy state in the Si band gap; however, thisAl adsorption mechanism is effectively used to form an Al-doped SiOx layer for this research aim.
Hence, in this experiment, Al precursors are intentionally added into the process solution.
Moreover, in order to form SiOx layers, an oxidant is required in the process solution. As
mentioned above, HNO3 can oxidize the Si surface and form ultrathin SiOx layers. However, insuch a strong acid, Al is not expected to adsorb on the Si substrate because metal contaminants,
including Al, exist as metal ions (cation), and they are repelled from the Si surface due to positive
zeta potential, dissolving into the acid solution. Therefore, one of the suitable Al precursors isaluminum nitrate (Al(NO3)3) because the material has a moderate pH of 5−6 with an oxidant
generated from nitrate (NO3−). Assumed by the proposed oxidation model by Asuha et al. [15]and M. Takahashi et al. [16], the nitrate ions are decomposed into O− ions at the surfaces with
SiO2 pore structures and the O− ions diffuse into SiO2 layers by electric field enhanced by theadsorbed NO3−, oxidizing Si surfaces by the diffused O− ions. Hence, the method can be simplified
by using Al(NO3)3 aqueous solution: just Si substrates are immersed into Al(NO3)3 aqueous
solution, forming simultaneously Al−O and Si−O in a low-temperature as an ultrathin Al-doped
SiOx layer. The remarkable feature is that this process is a simple wet chemical low-temperature
method, expecting high throughput and high-efficiency c-Si solar cells.

55

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HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

4.3 Film properties of Al-doped SiOx
4.3.1 Sample fabrication
In this research, double-side-polished 280 µm thick float-zone (FZ) n-type Si (100) substrates
with a resistivity of 2.5 Ω cm (TOPSIL PV-FZ™) were used. The Si substrates were cleaved into
20×20 mm-sized test sample pieces, followed by Radio Corporation of America (RCA) cleaning[17]. After 1% diluted hydrofluoric acid (HF) treatment for 1 min and rinse in deionized water(DIW) (resistivity of 18.2 MΩ cm) for 1 min, the hydrophobic Si substrates were immersed in
Al(NO3)3 aqueous solution with an Al(NO3)3·9H2O molar concentration of 0.25 mol L−1 for 5 to
30 min, as depicted in Figure 4.1. The hot Al(NO3)3 aqueous solution was prepared in advance by
dissolving Al(NO3)3·9H2O (17.8 g, special grade, Junsei Chemical Co., Ltd.) in DIW (190 mL)
in a Pyrex glass beaker (Corning Pyrex 600 mL) and heating to a boiling temperature (~95 °C)
on a hot plate (AS ONE Corporation, CHPS-170DF). During the heating process, the beaker was
capped to avoid the evaporation of the solution. The solutions were not intentionally stirred during
the process; however, drastically generated bubbles could be confirmed at the boiling temperature.
After the Al(NO3)3 treatment, processed Si substrates were rinsed in DIW for 1 min andcompletely dried by N2 blow.

Fig. 4.1 Schematics of the appearance for Al(NO3)3 oxidation treatment.

4.3.2 Sample evaluation
XPS (Shimadzu Corporation, Kratos AXIS-ULTRA DLD) measurement acquired Si 2p, Al 2p,
and O 1s core-level spectra, in which monochromatized Al Kα radiation (hν = 1486.6 eV) wasused for X-ray source. The generated photoelectrons were collected in the surface-normal
direction. The pass energy of photoelectrons was set to 20 eV. To calibrate all spectra, measured
Si 2p3/2 peaks were fixed to a binding energy of 99.2 eV [18]. The Al 2p and O 1s spectra werealso calibrated using these shifted binding energies of each sample. The backgrounds ofphotoelectrons were subtracted using the Shirley–Proctor–Sherwood method [19, 20], and the56CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

experimental spectra were peak fitted using Voigt functions. The fraction of Lorentzian
contributions and full width at half maximum (FWHM) of the peaks were not fixed to a particularvalue, and all peaks were optimized until the dissociations between experimental spectra andfitted curves were minimized using XPSPEAK 4.1 software. In this study, Si0 2p3/2, Si0 2p1/2, andSi4+ (SiO2) peaks were taken into account for the Si 2p peak fitting. Si0 2p1/2 area intensity was
fixed to the half of Si0 2p3/2 peak area intensity, and the splitting energy between Si0 2p3/2 and Si0
2p1/2 was set to 0.6 eV [18]. Each SiOx thickness was calculated using area intensity, atomic
density, photoionization cross-section [21], and the escape depths for Si and SiO2 [22],respectively. For quantitative analysis of the oxide, surface atomic concentrations of oxidized Al
and Si were calculated using atomic sensitivity factors of 0.283 for Si 2p, 0.711 for O 1s, and
0.193 for Al 2p [23]. The deconvolution of Si 2p core-level spectra was performed after fixing
the chemical shift for the suboxide components (Si1+, Si2+, Si3+) and the FWHM ratio of eachcomponent concerning the Si0. 0.95 [21], 1.75 [21], and 2.48 eV [21] were set as each chemical
shift and 1.57 [21], 2.07 [21], and 2.36 [21] as each FWHM ratio, respectively. In addition to that,
Si 2p and Al 2p core-level spectra were acquired with varying the take-off angle (TOA) of
photoelectrons in the range of 10 to 90° to investigate by using the surface sensitive conditions.
To confirm the Al-doping distribution semi-quantitatively, the oxidized Al and Si peak intensity
ratio was evaluated with respect to the varied TOA. FTIR measurements were conducted toanalyze the atomic density of the interfacial SiOx layers. The cross-sectional structure and thecomposition of the interface were evaluated using TEM-EDS.

4.3.3

Chemical bond states

Figure 4.2 shows the Si 2p and Al 2p core-level XPS spectra for samples treated in boiled
Al(NO3)3 aqueous solution with a molar concentration of 0.25 mol L−1 for various treatment times
in the range of 5 to 30 min. The measurement observed bulk-Si and SiO2-derived Si 2p peaks at
the binding energies of 99.2 and 102.5 eV, respectively. The SiOx intensities slightly increase withthe passing of treatment time. The binding energies of Al oxide can be assigned to be generally
larger than those of metal Al of 72.5–73.0 eV due to creating bonds with high electronegativity
element of O. [18, 24–26] For instance, the reported values are listed as follows: 75.8–76.6 eV[24]for amorphous and crystalline Al oxide, 74.7 eV[25] for γ-Al2O3, and 73.9 eV[26] for α- Al2O3.
Therefore, the observed peaks ~74–76 eV can be assigned as the Al–O bond states. No peak at~73 eV indicates the films do not contain metallic Al–Al states.

57

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HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

(a)

(b)
Si 2p

Al 2p

Fig. 4.2 (a) Si 2p and (b) Al 2p core-level X-ray photoelectron spectroscopy (XPS) spectra forsamples treated in boiled Al(NO3)3 aqueous solution with a molar concentration of 0.25 mol L−1
for various treatment times in the range of 5 to 30 min.

4.3.4

Atomic concentration and thickness

Figure 4.3 displays the computed atomic concentrations of oxidized Al and Si and the SiOx
thickness on samples treated in boiled Al(NO3)3 aqueous solution with a molar concentration of
0.25 mol L−1 for various treatment times in the range of 5 to 30 min. Although Si oxide peaks
gradually increase from 8.0 to 10.7 at.% with passing the Al(NO3)3-treatment time, the atomic
concentrations of all Al–O peaks are saturated in a range of 0.7–1.2 at.%. The estimated SiOx
thickness gradually increased from ~0.7 to ~1.2 nm. These results indicate that the growth of Al
oxides is finished in the initial stage, whereas only Si oxides can grow further. As mentioned in
4.2, nitrate ions are decomposed into O− ions at the SiO2 surfaces and diffuse into SiO2 layers,
and the migrated O− ions oxidize Si at the interface. Considering this O− diffusion oxidation modeland the growth rate difference between Al oxide and Si oxide, the fact implies that the Al-doped
SiOx layers consist of AlOy/SiOx/Si stack structures.

58

CHAPTER 4

Fig. 4.3

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEComputed atomic concentrations of oxidized Al and Si, and the SiOx thickness on

samples treated in boiled Al(NO3)3 aqueous solution with a molar concentration of 0.25 mol L−1
as a function of Al(NO3)3 treatment time.

4.3.5

Composition of sub-oxide and SiO2 component

Figure 4.4 illustrates the composition of sub-oxide (Si1+, Si2+, Si3+) and SiO2 (Si4+) components
obtained from deconvoluted Si 2p spectra on samples treated in boiled Al(NO3)3 aqueous solution
with a molar concentration of 0.25 mol L−1 for various treatment times in the range of 5 to 30 min.
The Si4+ component increases; on the other hand, the Si1+ component decreases with increasing
Al(NO3)3-treatment time from 5 to 30 min. D. Hiller et al. proposed that the tetrahedral structureof O-coordinated Al oxide is generated because the adjacent SiO2 imprints its tetrahedral nearorder and bond length into the Al–O MLs. [27] Considering this experimental fact, moretetrahedral structures of O-coordinated Al oxide favorably form with increasing Al(NO3)3treatment time.

59

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEComposition of sub-oxide (Si1+, Si2+, Si3+) and SiO2 (Si4+) components on samples

Fig. 4.4

treated in boiled Al(NO3)3 aqueous solution with molar concentration of 0.25 mol L−1 as a functionof Al(NO3)3 treatment time.

4.3.6

Al composition evaluated by surface sensitive condition

Figure 4.5 shows the Si 2p and Al 2p core-level XPS spectra for samples treated in boiled
Al(NO3)3 aqueous solution with molar concentration of 0.25 mol L−1 for 12.5 min acquiring withvarying the TOA of photoelectron in the range of 10 to 90° to investigate in the surface sensitive
conditions. As shown in Si 2p and Al 2p spectra, peak intensities gradually decrease withdecreasing TOA. To compare those spectra, using each atomic concentration of oxidized Al(XAlOx) and Si (XSiOx), each oxidized Al atomic concentration ratio (rAlOx), as described below, wasestimated.

𝑟AlOx =𝑋AlOx
𝑋SiOx + 𝑋AlOx(4.1)

Figure 4.6 describes the Al composition in the oxide layer as a function of TOA. Al composition
observed at TOA of 30° was minimal but linearly increased from 30 to 10°. Since surface sensitive
condition (small number of TOA) observes higher Al composition compared to interface sensitive60

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

condition (large number of TOA), Al oxides possibly exist at the surface of the SiOx layer. Moresurface-sensitive TOA conditions might show higher Al composition, but less than 10° of TOA
was not able to be set due to the limitation of the XPS setting. Hence, this result also indicatesthat the ultrathin Al-doped SiOx consists of a bi-layer of AlOy/SiOx on a Si substrate.

(a)Fig. 4.5

(b)

Si 2p

Al 2p

(a) Si 2p and (b) Al 2p core-level XPS spectra acquired with varying the TOA ofphotoelectron in the range of 10 to 90°.

Surface sensitiveFig. 4.6 Oxidized Al atomic concentration ratio as a function of TOA.
61

CHAPTER 4

4.3.7

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEAtomic density of interfacial SiOx

To deduce the atomic density of the interfacial SiOx layer, FTIR measurements were performedon Al(NO3)3-treated samples. For such an ultrathin layer, the off-plane vibration cannot be
detected by a standard setup of vertical incident light. In order to detect the longitudinal optical(LO) phonon vibrational mode of SiOx, the measurements were performed at an oblique incident
light setting of 74° by using a variable angular holder, which can be detected in LO mode by ppolarized light due to the Berreman effect [28]. Figure 4.7 displays FTIR spectra of samples
treated in boiled Al(NO3)3 aqueous solution with a molar concentration of 0.25 mol L−1 for various
treatment times. The transversal optical (TO) phonon modes were detected and confirmed nosignificant shift in the wavenumber range of 1029–1032 cm−1. On the other hand, significant shiftsof LO modes were confirmed and the LO peaks were shifted to the higher wavenumber from 1071
to 1207 cm−1 with increasing Al(NO3)3 treatment time. From the wavenumber of LO and TO
phonons modes, SiOx atomic densities (ρ) were calculated from the equation [29] below.
2
2
𝜈𝐿𝑂
− 𝜈𝑇𝑂
= 𝐶𝜌

(4.2)

Where C is a constant, and νLO and νTO are the frequency of LO and TO modes, respectively. As
the value of C, the reported value of 1.71×10−17 cm from bulk SiO2 [29, 30] was used. Figure 4.8
represents the deduced SiOx atomic densities as a function of Al(NO3)3 treatment time. This result
indicates that atomic density increases from 1.79×1022 to 2.24×1022 cm−3 with increasing Al(NO3)3
treatment time from 5 to 30 min. Hence, this atomic density increase agrees with the SiO2
composition increase with increasing treatment time, as described in Figure 4.4. The Al-doped
SiOx layer can be of higher quality, which is denser and has a lower concentration of sub-oxides.

62

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDELO

TOFig. 4.7 FTIR spectra of samples treated in boiled Al(NO3)3 (aq.) with a molar concentration of
0.25 mol L−1 for various treatment times.

Fig. 4.8

Deduced SiOx atomic density of Al-doped SiOx formed in boiled Al(NO3)3 (aq.) with amolar concentration of 0.25 mol L−1 for various treatment times.

63

CHAPTER 4

4.3.8

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDECross-sectional structure

Figure 4.9 illustrates the cross-sectional TEM image of the Al-doped SiOx/Si interface for samples
treated in boiled Al(NO3)3 aqueous solution with a molar concentration of 0.25 mol L−1 for 12.5
min, and Figure 4.10 shows EDS mapping concerning the depth: (a) O Kα, (b) Si Kα (c) Al Kα,
respectively. The thickness of the layer was determined to be 0.93 nm, and it is also recognizablefrom the detected TEM-EDS O Kα signals, as shown in Figure 4.10 (b). The thickness is estimatedto be close to 0.92 nm from XPS results. The important thing here is that a microscopic analysis
also revealed that the Al-doped SiOx layer can be recognized as an ultrathin layer. Moreover, AlKα detection was also performed to confirm the Al depth distribution, but the Al Kα signal was anoise level, as shown in Figure 4.10 (c).

Al-doped SiOx
Pd-Pt-C

Fig. 4.9

Si

Cross-sectional TEM image of Al-doped SiOx/Si interface for samples treated in boiled

Al(NO3)3 aqueous solution with a molar concentration of 0.25 mol L−1 for 12.5 min.

64

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

(a)

(b)

(c)

Fig. 4.10

Cross-sectional TEM-EDS mapping of Al-doped SiOx/Si interface in the same

viewpoint of the Cross-sectional TEM-image shown in Fig. 4.9: (a) O Kα, (b) Si Kα (c) Al Kα.
65

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HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

4.4 Surface passivation quality by Al-doped SiOx
4.4.1 Sample fabrication
Double-side-polished 280 µm thick FZ n-type Si (100) with a resistivity of 2.5 Ω cm (TOPSIL
PV-FZ™) were used in this study. The Si substrates were cleaved into 50×50 and 22×22 mmsized test sample pieces, followed by RCA cleaning [17]. After 1% diluted HF treatment for 1
min and rinse in DIW (resistivity of 18.2 MΩ cm) for 1 min, the hydrophobic Si substrates were
immersed in Al(NO3)3 aqueous solution with Al(NO3)3·9H2O molar concentration of 0.15, 0.25,
or 0.50 mol L−1 for 3 to 30 min. The hot Al(NO3)3 aqueous solution was prepared in advance by
dissolving Al(NO3)3·9H2O (10.7, 17.8, or 35.6 g, special grade, Junsei Chemical Co., Ltd.) in
DIW (190 mL) in a Pyrex glass beaker (Corning Pyrex 600 mL) and heating to a boilingtemperature (~95 °C) on a hot plate (AS ONE Corporation, CHPS-170DF). During the heatingprocess, the beaker was capped to avoid the evaporation of the solution. The solutions were notintentionally stirred during the process; however, drastically generated bubbles could be
confirmed at the boiling temperature. After the Al(NO3)3 treatment, processed Si substrates wererinsed in DIW for 1 min and completely dried by N2 blow. In addition, the process temperatureand duration of the Al(NO3)3 treatment were also varied from room temperature to boilingtemperature with an Al(NO3)3·9H2O molar concentration of 0.15 mol L−1 for 10 to 90 min.

4.4.2 Sample evaluationτeff measurements were performed using microwave photoconductivity decay (μ-PCD) [31](Kobelco Research Institute, Inc., LTA-1510EP). The τeff mapping data were acquired within thearea of 20×20 mm2 in 22×22 mm2 sized samples with a 0.5 mm pitch. A pulsed laser with awavelength of 904 nm and an areal photon density of 5×1013 cm−2 was used for the carrier
excitation. In addition, τeff and iVoc as a function of Δn were analyzed using QSSPC [32] (Sinton
Instruments, WCT-120) for 50×50 mm2 sized samples treated under one of the best conditions.
The relationship between τeff and Seff was used for finding Seff max values, which are described as[33]
τeff−1 = τbulk−1 + 2Seff W−1,

(4.3)

where τbulk and W represent the minority carrier lifetime in bulk c-Si and the thickness of Si
substrates, respectively. In this study, assumed τbulk = ∞, due to the exceptionally high τbulk of the66

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

Si substrates, Seff, max is thus defined asτeff−1 = 2Seff,max W−1(4.4)

and Seff,max was calculated using τeff at Δn =1×1015 cm−3, and W = 280 μm for wafer thickness. J0
value was also analyzed by the Kane and Swanson method [34] using the auger-corrected
effective minority carrier lifetime (τcorr) over the Δn range of 1×1015 to 9×1015 cm−3.

4.4.3 Surface passivation quality
Figure 4.11 represents the τeff and the corresponding maximum effective surface recombination
velocity (Seff, max) as a function of Al(NO3)3 treatment time for the various Al(NO3)3 molar
concentrations. Seff, max decreases to <26 cm s−1 and rapidly decreases with increasing Al(NO3)3
molar concentration from 0.15 to 0.50 mol L−1. The treatment time acquiring the minimum Seff,max value can be reduced from 17.5 to 10 min by the higher concentration of the Al(NO3)3 aqueous

solution. However, further longer treatment causes gradual degradation of surface passivation
quality, as shown in Seff, max increase. One of the τeff mapping data is depicted in Figure 4.12.
Except for the lower center region, the areal homogeneity of the τeff is good. Those τeff wereacquired using the typical excitation condition of 904 nm excitation laser light with areal densityτeff (µs)

of 5.0×1013 cm−2.

Fig. 4.11

τeff and the corresponding Seff, max as a function of Al(NO3)3 treatment time for thevarious Al(NO3)3 molar concentrations.
67

CHAPTER 4

Fig. 4.12

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEτeff mapping of the sample treated in boiled 0.25 mol L−1 Al(NO3)3 (aq.) for 12.5 min.

Figure 4.13 displays τeff and the corresponding Seff, max as a function of Al(NO3)3 treatment time
for the Al(NO3)3 molar concentrations 0.15 mol L−1 at various treatment temperatures from room
temperature to boiling temperature. As shown in the figure, τeff significantly increases withincreasing process temperature in a short treatment time. In addition to that, each temperature
condition has maximum τeff values and further longer treatment degrades the surface passivation
quality. Moreover, the maximum surface passivation quality can be obtained by increasing theprocess temperature relatively quickly. Besides τeff of 502 µs was achieved in a relatively shorttime (17.5 min) at boiling temperature, it took a longer time (240 min) to obtain a τeff of 271 µs atroom temperature. The possible reason for the existence of the maximum τeff value is discussed
below. Considering the induction of negative Qf originated from AlO4 structures formed bytrapping the excited electrons into Al-induced acceptor states in the Al-doped SiOx layer, more
adjacent to stoichiometric SiO2 stabilizes to create this AlO4 structure. As shown in XPS and FTIR
results, the SiO2 component is increased with increasing Al(NO3)3 oxidation treatment time.
However, excess treatment time make the further oxide grow at the interface, which suppresses
the electron tunneling through the oxide and trapping into the Al-induced acceptor states. Hence,
excess oxide growth inhibits the negative Qf induction, causing the τeff degradation.

68

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEτeff (µs)

CHAPTER 4

Fig. 4.13

τeff and the corresponding Seff, max as a function of Al(NO3)3 treatment time for the

Al(NO3)3 molar concentrations of 0.15 mol L−1at various treatment temperatures.
For one of the best Al(NO3)3-treatment conditions, a molar concentration of 0.25 mol L−1 for 12.5
min, QSSPC measurement was performed and confirmed a high τeff of 886 μs at Δn =1×1015 cm−3,
as shown in Figure 4.14. In particular, since the low Δn region ~1014 cm−3 shows high surface
passivation quality as τeff of >1 ms, the field effect might be related to such good surface
passivation, which can induce the band bending of the Si surface. The corresponding Seff, max was
estimated to be <16 cm s−1, and an iVoc = 670 mV at 1 sun illumination and J0 = 65 fA cm−2 wereτeff (ms)analyzed, as shown in Figure 4.15 and Figure 4.16, respectively.

Fig. 4.14

Δn dependence on τeff and Seff, max of the sample treated in boiled 0.25 mol L−1

Al(NO3)3 (aq.) for 12.5 min.
69

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

1/τeff −1/τAuger (s−1)

CHAPTER 4

Fig. 4.15

J0 = 65 fA cm−2

Δn dependence on Auger-corrected inverse effective lifetimes of the sample treated inboiled 0.25 mol L−1 Al(NO3)3 (aq.) for 12.5 min.

Fig. 4.16

Light intensity dependence on iVoc of the sample treated in boiled 0.25 mol L−1

Al(NO3)3 (aq.) for 12.5 min.

70

CHAPTER 4

4.5

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEElectrical properties of Al-doped SiOx/Si interface

4.5.1 Tunneling current density
Double-side-polished 280 µm thick FZ n-type Si (100) substrates with a resistivity of 2.5 Ω cm
(TOPSIL PV-FZ™) were cleaved into 20×20 mm-sized test sample pieces, followed by
performing RCA cleaning [17]. After 1% diluted HF treatment for 1 min and rinse in DIW(resistivity of 18.2 MΩ cm) for 1 min, the hydrophobic Si substrates were immersed in Al(NO3)3
aqueous solution with Al(NO3)3·9H2O molar concentration of 0.25 mol L−1 for 5 to 30 min. The
hot Al(NO3)3 aqueous solution was prepared in advance by dissolving Al(NO3)3·9H2O (17.8 g,
special grade, Junsei Chemical Co., Ltd.) in DIW (190 mL) in a Pyrex glass beaker (Corning
Pyrex 600 mL) and heating to a boiling temperature (~95 °C) on a hot plate (AS ONE Corporation,
CHPS-170DF). During the heating process, the beaker was capped to avoid the evaporation of
the solution. The solutions were not intentionally stirred during the process; however, drasticallygenerated bubbles could be confirmed at the boiling temperature. After the Al(NO3)3 treatment,
processed Si substrates were rinsed in DIW for 1 min and completely dried by N2 blow. After thelayer formation, Al circular electrodes were deposited on the Al(NO3)3-treated samples through amask by thermal evaporation in a vacuum chamber at a pressure of 9×10−4 Pa. Their areas were
estimated to be 0.0028 cm2 on average by observing the deposited electrodes using differential
interference contrast microscopy (Olympus Corporation, BX51), as shown in Figure 4.17 (a). The
deposited sides were covered with polyimide films, and the samples were immersed in 1% diluted
HF solution for 1 min to remove the oxide layer on the rear side of the samples. After the DIW
rinse, Al films were deposited on the rear sides of the HF-treated surfaces, which were scratchedwith a diamond pencil to form good ohmic contacts between Al electrodes and Si substrates. After
peeling off the polyimide films, J–V characteristics were evaluated using a precision
semiconductor parameter analyzer (Hewlett-Packard Company, 4156 A) mounted with a kelvintype four-terminal connection system. The J–V characteristics evaluated tunneling current density
for the metal–oxide–semiconductor (MOS) structure, as depicted in Figure 4.17 (b).

(a)

Fig. 4.17(b)(a) Deposited and used Al circular electrode and (b) metal–oxide–semiconductor(MOS) structure of Al circular electrode/Al-doped SiOx/n-Si(100)/Al.
71

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Figure 4.18 displays the J–V characteristics of MOS structures formed using the samples
treated in boiled 0.25 mol L−1 Al(NO3)3 aqueous solution for various treatment times. A significant
increase in leakage current density can be seen with the increasing treatment time when thenegative voltage is applied despite an extremely low voltage, especially in the samples treated for
10 min or further treatment time. This result implies that hole tunneling current density increases
with increasing Al(NO3)3 treatment time, originating from upward band bending of the Si surface
induced by negative fixed charges existing in the Al-doped SiOx layer, accumulating large
amounts of holes in the space charge region as mirrored charges of the negative Qf. In addition,
the formation of Al-induced acceptor states in the oxide layer might be increased with increasing
Al(NO3)3 treatment time, which might be related to the hole tunneling through the oxide layer.
Assuming the model that has been proposed, hole tunneling occurs via the Al-induced acceptor
states; the increase of tunneling current density is also attributed to the increase of the Al-induced
acceptor states with increasing the Al(NO3)3 treatment time. In such a circumstance, even though
the extremely low negative voltage is applied, the leakage current density is significantly
increased by the accumulated large number of holes at the Si surface. Therefore, the 5 min
Al(NO3)3-treated sample still shows low leakage current density due to the insufficient upward
band bending and formation of Al-induced acceptor states. The crucial evidence for the existenceof negative fixed charges in the oxide layer is demonstrated later in Section 4.5.4.

Fig. 4.18 J–V characteristics of MOS structures formed using the samples treated in boiled 0.25
mol L−1 Al(NO3)3 aqueous solution for various treatment times.

Figure 4.19 shows the leakage current density of the MOS samples at various applied negative
voltages concerning the Al(NO3)3-treatment time. The leakage current density at voltages of <0.1
V enlarges despite increasing SiOx thickness. This result implies that the Al-doped SiOx layers
72CHAPTER 4

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can realize both good surface passivation quality and sufficient hole tunneling and overcome theproblem of the conventional SiOx layers. More AlO4 structure is created in the SiO2 structure by
the longer Al(NO3)3-treatment time due to the structurally stable tetrahedral compounds. As
described in Figure 4.4, since the SiO2 (Si4+) component is gradually increased with increasing
Al(NO3)3-treatment time, the incorporated Al is possibly more stable as a formation of AlO4
structure, resulting in more formation of Al-induced acceptor states and the induction of upward
band bending of Si surface. In contrast, the AlO4 structure is hardly generated in the SiOx, withmany suboxides formed by the shorter Al(NO3)3-treatment time due to the structurally unstablecompounds. Considering that the trap-assisted tunneling current proportionally enlarges withincreasing the density of trap states and the depth of the level in the trap [35], the increased factors
might overcompensate the growing barrier thickness with increasing Al(NO3)3-treatment time,realizing an increase in a hole tunneling. Moreover, according to Figure 4.18, a current density of~40 mA cm−2 at a maximum power point for typical high-efficiency c-Si solar cells leads to a
voltage drop of −0.02 V for the samples with Al-doped SiOx formed by Al(NO3)3 treatment. This
voltage drop is much less than the reported value of 0.06 V for the samples with Al-doped SiOx
formed by ALD and rapid thermal oxidation (RTO) in the literature [4].

Fig. 4.19 Leakage current density of MOS samples with Al-doped SiOx formed in boiled 0.25
mol L−1 Al(NO3)3 (aq.) at various applied negative voltages concerning the Al(NO3)3-treatment
time.

73

CHAPTER 4

4.5.2

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEContact resistivity

For the evaluation of ρc, palladium (Pd) circular electrodes with various sizes (diameter of 0.3–
3.2 mm) were deposited on the surface of Al(NO3)3-treated p-type Si(100) substrates through amask by thermal evaporation in a vacuum chamber at a pressure of 9×10−4 Pa. For the rear sideof the samples, the surface was treated with a cotton swab dipped in 5% diluted HF aqueous
solution in advance until the hydrophobic surface was visually confirmed. The full-area Al
electrodes were thermally evaporated on the rear surfaces, which were scratched in advance toobtain a good ohmic contact. The Pd circular electrodes (Cox and Strack pattern) and metal–
oxide–semiconductor (MOS) structure of Pd circular electrode/Al-doped SiOx/p-Si(100)/Al are
shown in Figure 4.20.
Current–voltage (I–V) measurements were performed for each-sized electrode using a precision
semiconductor parameter analyzer (Hewlett–Packard Company, 4156 A) mounted with a kelvintype 4-terminal connection system in the dark at RT. I–V characteristics were measured when thenegative bias was applied to the Pd top electrode to acquire the hole tunneling current at the holeaccumulation condition for the p-type Si substrates. By the Cox and Strack method [36], Rtotal foreach electrode was calculated under negative voltages from ∼1 to ∼15 mV, corresponding to a
current density of ∼40 mA cm−2 at a typical maximum power point of c-Si solar cells [37]. Rspread
was calculated using the equation proposed in the literature [36], and each Rtotal − Rspread wasplotted concerning the inverse values of electrode areas. ρc was estimated by extracting the slope
obtained from the linearly fitted curve.

(a)

(b)

Fig. 4.20 (a) Deposited and used Pd circular electrodes (Cox and Strack pattern) and (b) metal–
oxide–semiconductor (MOS) structure of Pd circular electrodes/Al-doped SiOx/p-Si(100)/Al
referred to the structure in the literature [37].

Figure 4.21 (a) displays dark I–V curves obtained for circular electrodes of various sizes andFigure 4.21 (b) describes the Rtotal, calculated from Figure 4.21 (a), subtracted by Rspread (Rtotal −Rspread) as a function of inverse electrode areas. The fitted curve in Figure 4.21 (b) shows good
linearity and the ρc was calculated to be 20 mΩ cm2 from the slope of the fitted curve. The
74

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estimated value is one order lower magnitude compared with those of the ALD Al–O (1 ALD
cycle)/thermal-SiO2 stacks reported in the literature [37](a)Fig. 4.21(b)(a) Dark I–V curves obtained for circular electrodes of various sizes. The diametersof the electrodes (d) are shown in the inset. (b) Rtotal subtracted by Rspread (Rtotal − Rspread) as afunction of inverse electrode areas.

4.5.3 Sheet resistance
The sheet resistance (Rsheet) of Al-doped SiOx/n-Si(100) was measured using the four-point probe
method (Mitsubishi Chemical Analytech Co., Ltd., Loresta-AX MCP-T370) in the dark condition.
As a reference, the sheet resistance of non-doped SiOx/n-Si(100) and HF-treated n-Si(100) werealso evaluated. The non-doped SiOx with a thickness of ~1 nm was formed by immersing a Si
substrate into 70 wt.% concentrated nitric acid aqueous solution at room temperature for 10 min.
Besides, Rsheet was not able to be measured for the non-doped SiOx surface because the resistances
are higher than the limitation of the order of 107 Ω sq−1; Al-doped SiOx allowed measuring the
Rsheet and 5.85×105 Ω sq−1 was obtained on average of 18 measurement points in 30×25 mm sized
sample. This result, therefore, implies that holes are accumulated at the surface of Si, reducing
the Rsheet to a measurable level. However, considering the application for the solar cell, the value
is considerably huge as an emitter layer on the n-type c-Si solar cells. However, under the lightillumination condition, Rsheet was drastically reduced, two orders of magnitude lower than thosein the dark condition. In this experiment, the sample surface was illuminated at 1 sun intensity.
The Measurement setup of the Rsheet is displayed in Figure 4.22. Since the Al-doped SiOx layer isformed by a dipping process, the layer is formed on both sides of the sample. To avoid the shadowof the four-point probe, the measurements were performed on the rear side of the sample, whereas
the 1 sun was illuminated on the front side of the sample. This measurement can come into effect75

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

because of the sufficiently high bulk lifetime of the used Si substrate (τbulk >10 ms), which meansa longer diffusion length of hole in c-Si (~0.36 cm) compared to the thickness of the used Si
substrate (0.028 cm), leading to the sufficient hole accumulations at the surface of the rear side as
well. Eventually, an Rsheet of 4.86×103 Ω sq−1 was obtained under the 1 sun illumination. Thisdrastic reduction of Rsheet might be attributed to increasing the concentration of accumulated holes
at the Si surface due to further strong upward band bending induced by enhancement of trapping
excited electrons at the Al-induced acceptor states, existing in combination with Qf enlargement.
This result suggests that the inversion layer induced by an ultrathin Al-doped SiOx can possibly
be used as a p+ emitter for the n-type c-Si solar cells under light illumination. As a reference, anon-doped ultrathin SiOx layer with a thickness of ~1 nm was formed by immersing a Si substrate
into HNO3 aqueous solution for 10 min at RT. The sheet resistances in the dark and 1 sunillumination condition were not able to be measurable because of the high resistance over themeasurement limitation of >107 Ω sq−1. Hence, Al-doped SiOx demonstrated huge advantage withmore conductive layer less than two order magnitude of Rsheet compared to the non-doped SiOx
surface. The sheet resistances of Al-doped SiOx/n-Si(100) in the dark and under light illumination
are shown in Table 4.1.

Fig. 4.22 Measurement setup of the Rsheet under light illumination.

Table 4.1

Sheet resistances of Al-doped and non-doped SiOx/n-Si(100) in the dark and under

1 sun illumination.

Measurement condition

Al-doped SiOx/n-Si(100)Non-doped SiOx/n-Si(100)Dark

5.85×105 Ω sq−1

>107 Ω sq−1

Light illumination (1 sun)

4.86×103 Ω sq−1

>107 Ω sq−1

76

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4.5.4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

Determination of Fixed charge density and Interface state density

In general, to estimate Qf and Nit, capacitance–voltage (C–V) measurement is conducted andcalculated from the obtained capacitance of oxide (Cox) [38], but such an ultrathin layer hardly
evaluates these values due to the high leakage current, causing the lack of accuracy of measuring
Cox comes from an insufficient carrier accumulation at the interface of dielectric layer andsemiconductor. Thus, one of the promising solutions to analyze without such a lack of accuracy
is the estimation from Vbi generated by an ultrathin layer. As described in the literature [39],surface hole density (ps) is exponentially proportional to the Vbi, and Qf is proportional to thesquared ps. Assuming the dielectric layer with a sufficient negative Qf, the relationship among ps,
Vbi, and Qf is described as,V

𝑄𝑓 = − √2𝑘B 𝑇𝜀𝑠𝑖 pd exp( 𝑉bi ),𝑇Vbi

𝑝𝑠 = 𝑝𝑑 exp (𝑉 ),

(4.5)

(4.6)

𝑇

where pd is the concentration of holes at a depth of the flat band region and VT is the thermal
voltage, respectively. Hence, Qf can be easily estimated from the measured and analyzed Vbi value.
In addition, Nit is proportional to Seff and can be described as follows.

𝑁𝑖𝑡 =

𝑆eff
2𝑘 𝑇𝜀
𝜎n 𝜈n (𝑁dop +𝛥𝑛)( B 2 si )
𝑄

(4.7)

f

Where σn is the electron capture cross section, and νn is the velocity of the electron, respectively.
Hence, Nit can also be estimated from the analyzed Seff value from measured τeff.
For the determination of Qf and Nit, the oxide on the rear side of the Al(NO3)3-treated sample
was removed with a cotton swab dipped in 5% diluted HF aqueous solution in advance until thehydrophobic surface was visually confirmed. The full-area Al electrodes were thermallyevaporated on the rear surfaces, which were scratched in advance to obtain a good ohmic contact.
Subsequently, a C–V measurement was performed to obtain depletion capacitance in the negative
voltage region using a mercury (Hg) probe. The Vbi was evaluated by Mott-Schottky plot analysis[40], linearly fitted capacitance in the voltage range of −2.5 to 0 V.
Figure 4.23 depicts the Mott-Schottky plot for n-type Si(111) samples with Al-doped SiOx
77

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formed in boiled 0.25 mol L−1 Al(NO3)3 aqueous solution for 12.5 min. Vbi was extracted to be
0.870 V from the interception of the extrapolated fitted line concerning the voltage. The large Vbi
of Al-doped SiOx/n-Si means that strong upward band bending is induced at the surface of n-Si.
This is crucial evidence for the existence of a large number of negative charges in the Al-doped
SiOx layer, and Qf was calculated to be −3.21×1012 cm−2. Furthermore, assumed σn is 1.8×10−14
cm2, determined for Al2O3/Si in the literature [41], the Nit was estimated to be 1.43×1012 cm−2
using Seff, max of 23.3 cm s−1, which was analyzed from measured τeff of Al-doped SiOx/n-Si(111).
This Nit value is very close to the Nit of 1.03×1012 cm2 for HNO3 grown SiOx, ~4.03 × 1012 cm2for HCl grown SiOx, and ~2.73×1012 cm2 for H2SO4 + H2O2 grown SiOx [12]. Hence, the surface
passivation mechanism originated from chemical passivation and field-effect passivation.
However, in general, since the wet chemical oxide shows very low τeff of ~5 µs, more than twoorders of magnitude lower than Al-doped SiOx, the field effect passivation is the main scheme forthis surface passivation mechanism. The determined Qf and Nit are summarized in Table 4.2.

(a)

(b)

Fig. 4.23 (a) Mott-Schottky plot for the samples of n-type Si(111) with Al-doped SiOx formed
in boiled 0.25 mol L−1 Al(NO3)3 (aq.) for 12.5 min. (b) Metal–oxide–semiconductor (MOS)
structure of Hg probe/Al-doped SiOx/n-Si(111)/Al.

Table 4.2

Determined Qf and Nit

Sample structure

Qf (cm−2)

Nit (cm−2)

Al-doped SiOx/n-Si(111)−3.21 × 1012 cm−2

1.43× 1012 cm−2

Using PC1D simulation software, the estimated Qf values were confirmed to determine whether
they were reasonable. Assuming an estimated Qf value of −3.21 × 1012 cm−2 on the n-type Si
substrate with a Ndop of 2×1015 cm−3 and a thickness of 200 µm, the ps, surface electron density78CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE(ns) and Vbi were calculated to 3.974×1019 cm−2, 2.516 cm−2, and 0.887 V, respectively. Hence, thesimulated Vbi agrees well with the measured value of 0.870 V. The simulated carrier densities andenergy bands are shown in Figures 4.24 and 4.25. These simulated results show that the c-Si
surface is strongly inverted over the depth of ~250 nm in the short-circuit and under the dark
condition.

Electron

HoleFig. 4.24 Depth profile of carrier densities for the negatively charged n-type c-Si surface.

Ec
EF

Ev

Fig. 4.25 Depth profile of energy bands for the negatively charged n-type c-Si surface.

79CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

4.6 Carrier selectivity of Al-doped SiOx
Here, the selectivity (S10) [42] was estimated to be 13.3 by using the equation below with two
analyzed values of J0 = 65 fA cm−2 and ρc = 20 mΩ cm2,

𝑉

𝑆10 = log10 (𝐽 th
),
𝜌

(4.8)

0 c

where Vth represents thermal voltage. In addition, assuming the optimum contact fraction fmax = 1,
the maximum efficiency (ηmax) can be calculated using S10 obtained in the equation below [42],1𝜂max (𝑆10 ) = [(2.452 𝑆10 − 4.240)−19.52 + (29.21)−19.52 ]− 19.52 .

(4.9)

Since the ηmax can be calculated to 27.7% with S10 of 13.3, this carrier-selective contact has a highpotential because of a comparable passivating contact with other technology that has beenreported [37, 43–52]. Note that since Ag electrode degraded surface passivation as described in
Chapter 5, but the J0 was analyzed for a no Ag deposited sample and the ρc was analyzed using pSi as a substrate and Pd as a contact electrode, the analyzed J0 and ρc are not the practical value
due to the different sample structure from the real Al-COIL device structure (Ag/Al-doped SiOx/nSi). To compare other carrier-selective passivating contacts, the relationship among J0, ρc, and S10
with various carrier-selective passivating contacts are displayed in Figure 4.26. The red star
represents the carrier selectivity of Al-doped SiOx analyzed in this study. It can be recognized that
the S10 value is comparable to other technology, but it must be further improved to achieve highefficiency c-Si solar cells close to the theoretical limit.

Fig. 4.26 J0, ρc, and S10 for various passivating contacts. Each solid line shows S10 values from10 to 18.
80

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4.7

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

Effect of surface orientation

4.7.1 Surface passivation on Si(100) and Si(111)
Figure 4.27 shows τeff and the corresponding Seff, max as a function of Al(NO3)3 treatment time forthe various Al(NO3)3 molar concentrations from 0.15 to 0.50 mol L−1 for Si(100) and Si(111)substrates. It can be seen τeff improvement for Si(111) substrates relatively in a short time,
especially in a low Al(NO3)3 concentration condition, shows a significant difference in the process
time by reaching the maximum τeff. However, both orientation shows a significant increase in τeff
for a short process time with increasing the Al(NO3)3 molar concentration; it is presumably due
to the fast oxidation. For all concentration conditions, Si(100) enabled higher surface passivation.
For instance, at the condition of 0.50 mol L−1, 10 min treated (100)-oriented sample shows arelatively high τeff of 556 µs but 5 min treated (111)-oriented sample shows a lower τeff of 423 µs.
Besides, Si(111) showed a partially high τeff region can be observed for 7.5 min in 0.15 mol L−1,
as shown in Figure 4.28. Furthermore, the highest maximum τeff was 933 µs. Since such a shortprocess time allows us to obtain excellent τeff and a higher concentration process did not showsuch a high maximum τeff, Si(111) might have a narrow process window. The lower concentrationand longer treatment time will be a promising solution for obtaining a higher surface passivation
for the Si(111) surface. As observed in Figure 4.11, further Al(NO3)3 treatment caused the τeff
degradation for all conditions and both orientated substrates. This might have originated from alack of electron trapping in the Al-induced acceptor states due to the thickened SiOx formed by
the more prolonged treatment suppressing the electron tunneling through the SiOx layer from Si
to the acceptor states, resulting in low induction of the negative fixed charges and the suppressionof the upward band bending of Si surface. In particular, this significant degradation for Si(111)
was confirmed at the relatively high condition of 0.50 mol L−1. H. Angermann et al. reported that
Si(111) shows lower Nit compared with Si(100) in wet-chemical passivation schemes such as SiO2
formed by RCA or hot water. They proposed that a faster oxidation rate for Si(100) leads to higher
Nit, originating from lower stability of H-terminated Si(100) surfaces against reoxidation in air.
In addition, the electronegativity difference between the two bonded substituents such as H, OH,
and F and the Si occurs a stronger polarization of the Si back bond in Si(100), leading to a heavy
attack of oxidants [56]. However, in the case of Al(NO3)3 oxidation, the oxidation reaction on
Si(111) seems faster than on Si(100). If so, there is a discrepancy in the oxidation rate, but therelatively low surface passivation on Si(111) can be explained by this degradation model.

81

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEτeff (µs)(a)τeff (µs)(b)τeff (µs)(c)

Fig. 4.27

τeff and Seff,max as a function of Al(NO3)3 treatment time for the various Al(NO3)3 molar

concentrations: (a) 0.15, (b) 0.25, (c) 0.50 mol L−1 for Si(100) and Si(111) substrates.

82

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Fig. 4.28 τeff mapping of Al(NO3)3-treated Si(111) for 7.5 min in 0.15 mol L−1 Al(NO3)3 (aq.).

4.7.2 Chemical bond states on Si(100) and Si(111)
Figure 4.29 illustrates the Si 2p and Al 2p core-level XPS spectra for Si(111) and (100) samples
treated in boiled Al(NO3)3 aqueous solution with a molar concentration of 0.25 mol L−1 for various
treatment times in the range of 2.5 to 30 min. The observed peak around 102.5–103 eV and ~75
eV can be assigned to SiO2 (Si4+) and Al–O, respectively.
(b)

(a)

Si 2p

Al 2p

Fig. 4.29 (a) Si 2p and (b) Al 2p core-level XPS spectra for Si(111) and (100) samples treatedin boiled Al(NO3)3 (aq.) with a molar concentration of 0.25 mol L−1 for various treatment times.
83

CHAPTER 4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDETheir peaks were fitted with the Voigt function, and the evaluated SiO2 composition, SiOx
thickness, and oxidized Al atomic concentration as a function of Al(NO3)3 treatment time areshown in Figure 4.30 (a), (b), and (c), respectively. As expected, the SiOx on Si(111) was thicker
than that on Si(100) and formed in a short treatment time. Hence, Si(111) has a higher growth ratein the initial oxidation reaction. Depending on the higher growth rate, SiO2 composition is also
higher value from the initial oxidation reaction. However, the thickness and the SiO2 composition
for both orientations are almost the same at 30 min treatment time. Similarly, the oxidized Al
atomic concentration of Si(111) is higher than that of Si(100) in a short process time of 5 min.
After that, the atomic concentration slightly fluctuated between 1.2–1.4 at.% and saturated at ~1.2
at.% for both orientations. Hence, Al oxides are formed in the initial reaction stage, especially on
the Si(111) surface. Considering the fast oxide growth rate of Si(111), the faster oxidation ratecauses higher Nit on Si(111). One possible reason for the higher oxidation rate is that theterminated OH on the Si(111) has a higher density of electrons due to the larger number of Si
back bonds, providing electrons from three back bond Si to one terminated OH. Hence, a
nucleophilic Si-OH is highly reacting with Al precursor as a Lewis acid, and negatively charged
Al oxides might assist the O− diffusion to the Si surface by the repulsion force between AlO4− andO−, causing faster formation of Al oxides and SiOx.

(a)

(c)

(b)

Fig. 4.30 (a) SiO2 composition, (b) SiOx thickness, and (c) oxidized Al atomic concentration as afunction of Al(NO3)3 treatment time for Si(111) and (100) samples with Al-doped SiOx formed inboiled Al(NO3)3 (aq.) with a molar concentration of 0.25 mol L−1 for various treatment times.

84

CHAPTER 4

4.7.3

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDEElectrical properties on Si(100) and Si(111)Tunneling current density
Figure 4.31 represents J–V characteristics of MOS structures on Si(111) and Si(100) formed using
the samples treated in boiled 0.25 mol L−1 Al(NO3)3 aqueous solution for various treatment times
from 2.5 to 30 min in the applied voltage range of −1 to 0 V. The sample fabrication and evaluation
were performed using the same procedure and method as described in 4.5.1. For all the treatment
time conditions, Si(111) samples show higher leakage current density, meaning higher tunnelingcurrent density through the Al-doped SiOx on Si(111) substrates. In addition, the significant
difference was confirmed at the short treatment time, especially in 2.5 min, four orders ofmagnitude higher than that of Si(100) substrate. Furthermore, the leakage current density was
gradually increased with increasing treatment time, while the 30-minute treated sample shows anorder decrease in the leakage current density due to the weakened upward band bending, resulting
in a reduction in the hole selectivity. This result shows the same tendency with the amount of SiO2
component and SiOx thickness except for 30 min treated samples. The higher leakage current
density on Si(111) is highly related to the higher SiO2 content and larger SiOx thickness. Moreover,as shown in Figure 4.30 (b), in a short treatment time until 10 min, τeff of Si(111) showed higher
values. Hence, these results imply that Si(111) shows faster induction of hole selectivity becauseof the faster oxide growth rate.
(a)

(b)

(c)

(d)

(e)

Fig. 4.31 J–V characteristics of Al circular electrode/Al-doped SiOx/n-Si/Al MOS structures on
Si(111) and Si(100) formed using the samples treated in in boiled 0.25 mol L−1 Al(NO3)3 aqueous
solution for various treatment times: (a) 2.5 min, (b) 5min, (c) 7.5 min, (d) 12.5 min, and (e) 30
min.

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Sheet resistance on Si(100) and Si(111)
Figure 4.32 describes an Rsheet of Al-doped SiOx/n-Si(100) or (111) under the dark or 1 sunillumination. Si(111) shows one order of magnitude lower Rsheet compared to Si(100) in the dark
condition. Moreover, 1 sun illumination realized a further two orders of magnitude decrease in
Rsheet, obtaining 317 Ω sq−1 on average. Since solar cell device operates under 1 sun illumination
as a standard condition and the conventional c-Si solar cell has a pyramidal textured Si surface
with (111)-oriented facet, the result under the 1 sun illumination condition demonstrated practicalvalues as a comparable Rsheet, even though it is still one order higher than that of 50–85 Ω sq−1 forthe conventionally formed B-diffused p+ emitter surfaces [47]. Regarding tunneling current
density and sheet resistance, the Si(111) surface has a higher potential for hole-selective contacts.
In addition, as described above, since the pyramidal textured Si surface consists of (111)-facets,
it will be favorably implemented for this type of hole-selective contact.

Fig. 4.32 Rsheet of Al-doped SiOx/n-Si(100) or (111) in the dark or under 1 sun illumination.

86

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4.7.4

HOLE-SELECTIVE CONTACT BY AN ULTRATHIN AL-DOPED SINICON OXIDE

Effect of light illumination

Influence on sheet resistance
To investigate the influence of light illumination on the sheet resistance, Rsheet values wereacquired under the monochromatized light with a wavelength of 550 nm and various areal photon
densities in the range of 5×1012 –1×1014 cm−2. Figure 3.33 (a) shows a measuring sample underlight illumination with a wavelength of 550 nm and an areal photon density of 1×1014 cm−2. Figure
4.33 (b) shows Rsheet values of Al-doped SiOx/n-Si(111) measured under the monochromatizedlight illumination with a wavelength of 550 nm as a function of areal photon density. As a result,
Rsheets were slightly decreased with increasing light intensity. Since the obtained data plots can belinearly fitted with a good correlation coefficient of 0.996, Rsheet positively correlates to lightillumination intensity.

(a)

(b)

Fig. 4.33 (a) Measuring sample under the light illumination with a wavelength of 550 nm andan areal photon density of 1×1014 cm−2. (b) Rsheet of Al-doped SiOx/n-Si(111) measured under the
monochromatized light illumination with a wavelength of 550 nm as a function of areal photon
density.

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Influence on surface passivation
Figure 4.34 (a) shows τeff of Al(NO3)3-treated n-Si(111) in boiled 0.25 mol L−1 Al(NO3)3 aqueous
solution for 12.5 min as a function of areal photon density for an excitation pulsed laser light witha wavelength of 349 nm. The τeff were acquired using µ-PCD. As described in the figure, τeff wassignificantly increased with increasing areal photon density, and this effect cannot be observed
for the non-doped SiOx/n-Si structure. This increase might be attributed to the enhancement of
the negative Qf induced by electron trapping in the Al-induced acceptor states due to the instantly
generated excess carriers during the illumination of the pulsed laser light, increasing the upward
band bending and enhancement of field-effect passivation.
In addition, the τeff measurement under the steady state light illumination was performed. Figure
4.34 (b) shows τeff of Al(NO3)3-treated Si(111) in boiled 0.25 mol L−1 Al(NO3)3 (aq.) for 12.5 min
as a function of the intensity of steady-state light. A pulsed laser light with a wavelength of 904
nm was used to generate excess carriers. Generally, although τeff decreases under the steady state
light illumination due to reducing the band bending, resulting in the degradation of field-effectpassivation. However, in this Al-doped SiOx/n-Si case, a drastic increase of τeff was confirmed
under the illumination with a low-intensity steady state light below 0.7 suns. This increase mightalso originate from the enhancement of the negative Qf by the steady-state light illumination,
resulting in the upward band bending increase and enhancement of field-effect passivation.
(b)τeff (µs)τeff (µs)(a)

Fig. 4.34 (a) τeff as a function of areal photon density for an excision pulsed laser light with awavelength of 349 nm. (b) τeff as a function of the intensity of steady-state light measured using
an excision pulsed laser light with an areal photon density of 1×1013 cm−2 and wavelength of 904
nm. Both measurements were conducted for the same sample.
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Moreover, Δn dependency of τeff was investigated before and after light illumination. Figure
4.35 illustrates τeff as a function of Δn acquired using QSSPC before and after the illumination oflaser light with a wavelength of 349 nm and an areal photon density of 1.1×1015 cm−2, mounted
in a µ-PCD system. The laser light was illuminated on the Al(NO3)3-treated Si(111) sample in thearea of 44×44 mm with a 1 mm pitch. As clearly seen in Figure 4.35, after the laser lightillumination, τeff was significantly increased, especially in the low injection level of 1014 cm−3.
This result implies that activation is necessary to obtain high surface passivation, which can berealized by light illumination. Unoccupied Al-induced acceptor states capture the laser-excitedelectrons from the Si side, leading to highly negative Qf and field-effect passivation. The
significant increase at the low injection level suggests that the surface passivation comes from theτeff (µs)

field effect induced by the occupied Al-induced acceptor states.

Fig. 4.35

τeff as a function of excess carrier density (Δn) acquired using QSSPC before and after

the laser light illumination with a wavelength of 349 nm and an areal photon density of 1.1×1015
cm−2.

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4.8 Model of hole-selective passivating contact
Figure 4.36 (a) shows a model for the induction of negative Qf in the Al-doped SiOx layer on the Si
substrate after or under light illumination. As demonstrated in this experiment, light illumination
highly influences the surface passivation and hole tunneling properties. The negative fixed charges are
generated due to the captured excited electron in the Al-induced acceptor states existing in the Aldoped SiOx layer, occurring upward band bending of the Si surface and creating an inversion layer asa hole accumulation layer. Suppose the electrode is deposited on the Al-doped SiOx layer. In that case,the generated holes selectively tunnel through the oxide layer because the minority (electron) barrier
is quite high while the majority (hole) barrier is significantly reduced by the creation of Al-induced
acceptor states located below the valence band edge of Si [4]. Hence, electrons are repelled from thesurface due to the strong upward band bending induced by the field effect and the high minority barrier,while holes are selectively tunneling through via the Al-induced acceptor states, possibly occurring
trap-assisted tunneling; therefore, electrons and holes are effectively separated by the field-effect at
the surface of Si. This model for the hole-selective passivating contact is schematically described in
Figure 4.36 (b). In this configuration, the inversion layer possibly acts as a p+ emitter of the n-type cSi solar cells. Chapter 5 demonstrated this model as a novel device architecture fabricated by a simple
wet chemical process.

(a) Light illumination(b)

Fig. 4.36 (a) Model for the induction of negative Qf in the Al-doped SiOx layer on Si substrate
after or under light illumination. (b) Model for the hole-selective passivating contact of the Aldoped SiOx layer on Si substrate.

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4.9 Summary
This chapter developed an ultrathin Al-doped SiOx on c-Si substrates by a simple wet chemical
process using Al(NO3)3 aqueous solution. The excellent hole-selective passivating contact
properties (Seff of <16 cm s−1, iVoc=670 mV, S10=13.3) were confirmed despite the ~1 nm thickultrathin layer. The ultrathin layer consists of a bilayer of AlOy/SiOx with an Al concentration of~1 at.%. The C–V measurement and simplified extended SRH formalism revealed the ultrathin
layer has highly negative charges. The Qf value was estimated to be −3.2×1012 cm−2, contributing
excellent hole selectivity by induction of strong upward band bending. Since the estimated Nit
value was almost the same level as any other wet chemical oxide, the charge-induced field effectis dominant. Si(111) surface also showed excellent surface passivation, greater hole tunnelingproperty, and lower sheet resistance. Interestingly, after the light illumination, the surface
passivation quality was significantly increased, and sheet resistance was also significantly
improved under 1 sun illumination. The light-illuminated activation effect assists in elucidating
the mechanism for the hole-selective passivating contact in the ultrathin Al-doped SiOx. The
increased negative charges might originate from the captured excited electrons in the trap state
formed in the ultrathin layer.

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CHAPTER 5

AL-INDUCED CHARGED OXIDE INVERSION LAYER SOLAR CELLSCHAPTER 5
AL-INDUCED CHARGED OXIDEINVERSION LAYER SOLAR CELLS

The wet chemically grown ultrathin Al-doped SiOx provides surface passivation and holeselective contact at the same time. The negatively charged oxide induces an inversion layer in the
n-type c-Si surface, which originates from electron trapping in the Al-induced acceptor states
existing in the oxide as a formation of AlO4 structure. Surprisingly, since the excited electron
activates the hole selectivity in the Al-doped SiOx layer, the hole-selective passivating contact
properties are initiated under light illumination. In contrast, the surface passivation and theelectrical properties are inferior under dark conditions, as described in Chapter 4. Fortunately,
since solar cells certainly operate under light illumination, the light-activated ultrathin Al-doped
SiOx is expected to be able to be used under the practical 1 sun illumination condition. Hence,
expectedly, the light-induced hole-selectivity allows us to realize that the wet chemically grown
Al-doped SiOx acts as a p+ emitter for the n-type c-Si solar cells.
In this chapter, the concept of the novel Si-based solar cells is proposed, with an inversion layer
induced by the ultrathin Al-doped SiOx. I named it Al-induced charged oxide inversion layer (AlCOIL) solar cells, and the first fabricated Al-COIL solar cell was demonstrated following a simple
process and structure. This demonstration allows us to confirm the device's operation andelucidate the mechanism, as shown in the last chapter.

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5.1 Concept of Al-COIL architecture
The Al-COIL architecture conceptualizes to realize excellent surface passivation and carrier
selective contact using an ultrathin dielectric layer formed by a simple wet chemical process. The
recent study demonstrated excellent surface passivation of Seff <20 cm s−1 by a wet chemicallyprocessed ultrathin SiOx layer treated by rigorous cleaning processes and using superacidic(bis(trifluoromethane)sulfonimide crystals) solutions under inert conditions. In contrast, theAl(NO3)3 treatment also provides an excellent surface passivation of Seff, max <16 cm s−1 by AlCOIL architecture, and it can be recognized as a more industrially feasible process considering a
wet chemically grown ultrathin oxide layer with a thickness of ~1 nm formed by a simple wet
process using a relatively low-cost and safer chemical under atmospheric pressure conditions. In
addition, the surface passivation of the superacidic treated-ultrathin SiOx is mainly attributed to
the reduction of the Nit and minor contribution of the negative Qf of the order of 1010 cm−2 [1].
Hence, this surface passivation layer does not expect the hole-selective contact because thesurface hole density induced by a Qf of −4.5×1010 cm−2 [1] can be calculated to 2.65×106 cm−3from PC1D simulation, which cannot invert the Si surface. Therefore, to enable hole-selective
passivating contact through ultrathin SiOx formed by a simple wet chemical process, only AlCOIL is the promising solution, which can invert the n-type Si surface. Such an excellent ultrathin
passivation layer enables sufficient carrier tunneling without degrading surface passivation
quality, rather than hole-selective carrier tunneling due to the hole accumulation at the Si surface
as the negatively charged oxide layer induces the mirrored charges.
The fabrication process flow, the device structure, and the principle of the device operation are
shown in Figure 5.1. The Al-COIL solar cells are expected to reduce the processing steps with acost-effective, simple wet chemical process. In addition, since impurities such as B or Al are notdiffused into the Si surface, surface recombination can be mitigated from the recombination center
generated by the impurities. The following leading cell structure of TOPCon requires more
processing steps, and the cell efficiency is limited by Auger recombination in the front boron
emitter [2]. For the SHJ structure, the solar cell performance is limited by parasitic absorption in
the front a-Si passivation layer [3]. Thus, the inversion layer induced by a negatively charged
ultrathin SiOx layer can not only reduce the processing steps but also achieve higher efficiency
due to avoiding the Auger recombination and parasitic absorption loss in the emitter. However,
one of the biggest problems would be a difficulty to obtain an enough lateral conductivity throughthe inversion layer.

98CHAPTER 5

Fig. 5.1

AL-INDUCED CHARGED OXIDE INVERSION LAYER SOLAR CELLS

Fabrication process flow, the device structure, and the principle of the device operation.

The negative charge density must be sufficiently high to obtain conductive enough to becomparable to the B-diffused emitter. In the previous study, fixed charges in various oxide layers
have been investigated for the induction of an inversion layer in the Si surface [4, 5, 6]. However,this approach also finds it challenging to provide sufficiently high charge densities in the dielectric
layer, mitigating lateral resistive losses. Although, for a p-type base, a very high positive charge
density of 1.1×1013 cm−2 has been realized by silicon nitride (SiNx) [7], the charge density was
not high enough to introduce a sufficient lateral conductivity compared with conventional Pdiffused emitters. M. Yu et al. reported that the simulation revealed a positive charge density of
4×1013 cm−2 can provide a sheet resistance of the inversion layer as low as ~180 Ω sq−1 [8], which
is close to that of 130 Ω sq−1 for a P-diffused emitter [9]. Therefore, such a highly charged ultrathinoxide layer must be formed to accomplish carrier-selective contact with a comparable level oflateral conductivity on the front surface of other types of solar cells, such as PERC, TOPCon, andSHJ. As shown in Figure 1.3, only 6.7% charged Al in the Al-O monolayer induces such a highcharge density of −4×1013 cm−2 and very low Seff of 0.1 cm s−1 simultaneously. Moreover, PC1D
simulation suggests that the negative Qf shows a comparable sheet resistance of 119 Ω sq−1 to a
B-diffused emitter [9]. The Rsheet was calculated from the simulated hole density and hole mobility
concerning the depth using the equation below.

𝑅sheet =1𝑧
𝑞 ∫0 𝜇(𝑧) 𝑝(𝑧) 𝑑𝑧(5.1)

Where μ is hole mobility, p is hole density, and z is depth from the surface. The calculated Rsheet
concerning the Qf is shown in Figure 5.2.

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Fig. 5.2 Rsheet concerning the Qf calculated from PC1D simulation.

Hence, since Rsheet reduces inversely proportional to the Qf2, improving Qf is the most effective
approach to increase the lateral conductivity. However, considering achieving a 317 Ω sq−1 value
under 1 sun illumination, Qf might be enhanced to ~ −2×1013 cm−2, as following this relationshipbetween Rsheet and Qf. Therefore, the concept of this novel simple device will possibly be realized
as a high-efficiency, cost-effective c-Si solar cell.

5.2 Simulation for Al-COIL solar cells
The J–V characteristics and power densities were simulated using PC1D software. The device
structure with surface charge-induced emitter, which is assumed to be an Al-COIL and B-diffusedemitter, and these input parameters for PC1D simulation are shown in Figures 5.3 and 5.4,
respectively. For the surface charge, the front surface charge density was set to −3×1013 cm−2, andfor the B-diffused emitter, the typical B-doping concentration was set to 1×1019 cm−3 [10] with adepth factor of 1.135 µm which can be set a sheet resistance as 95 Ω sq−1 [10]. The simulations
were performed under the steady state condition of 1 sun illumination (steady state intensity of
0.1 W cm−2 with an AM 1.5 Global spectrum). In this simulation, the internal series resistance isassumed to be sufficiently low (1×10−6 Ω cm2). In addition, it is not considered an optical loss
attributed to the front reflectance but considered internal light diffusive reflection from the rear100CHAPTER 5

AL-INDUCED CHARGED OXIDE INVERSION LAYER SOLAR CELLS

surface. Figure 5.5 demonstrated excellent J–V characteristics even without a B-diffused emitter.
As shown in Table 5.1 of the simulated solar cell performance, η was calculated to be 27.21%.
Since the calculated η is very close to the theoretical limit of SHJ solar cells (27.5%) [11], it canbe recognized the surface charge-induced emitter device is expected to realize high-efficiency cSi solar cells. Figure 5.6 shows a PC1D simulated band energy diagram of the front surface at a
maximum power point. As shown in the diagram, the generated large quasi-fermi level splittingof 0.676 eV indicates that the negative charges from the surface induce an inversion layer as a p+emitter of n-type c-Si solar cells. The J–V characteristics for the conventional B-diffused device
and the simulated solar cell performance are shown in Figure 5.7 and Table 5.2, respectively.
Compared with the two types of solar cells, the surface-charged-induced device can be superior
to the conventional B-diffused device. Hence, these simulation results demonstrated that thesurface charge, in principle, induces a p+ emitter for the high-efficiency c-Si solar cell. Note that
PC1D simulation neglects the effect of sheet resistance. Hence, the simulated results display
values when the sheet resistance is ideal.

Fig. 5.3

Device structure and input parameters for PC1D simulation to obtain ideal J–V

characteristics of the device with the surface charge-induced emitter.
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Fig. 5.4

AL-INDUCED CHARGED OXIDE INVERSION LAYER SOLAR CELLS

Device structure and input parameters for PC1D simulation to obtain ideal J–V

characteristics of the device with B-diffused emitter.

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Fig. 5.5 PC1D simulated J–V curve and power density for surface charge-induced emitter.

Table 5.1

PC1D simulated solar cell performance for surface charge-induced emitter.
Jsc (mA cm−2)

Voc (mV)FF (%)

η (%)

41.50

0.746

87.91

27.21

Fig. 5.6 Band energy diagram of the front surface at the maximum power point for surface
charge-induced emitter.
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Fig. 5.7 PC1D simulated J–V curve and power density for B-diffused emitter.

Table 5.2

Fig. 5.8

PC1D simulated solar cell performance for B-diffused emitter.

Jsc (mA cm−2)

Voc (mV)FF (%)

η (%)

41.45

0.736

86.43

26.37

Band energy diagram of the front surface at maximum power point for B-diffusedemitter.
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Moreover, Figure. 5.9 shows the simulated Voc as a function of S for the device structure withsurface charge-induced emitter and B-diffused emitter. Significant Voc reduction was confirmedin the device structure with B-diffused emitter in the range of S >102 cm s−1 order. In comparison,such a significant reduction cannot appear by the order of 106 cm s−1 for the device structure withsurface charge-induced emitter. This is considered to be attributed to the significant difference inthe surface carrier densities between electrons and holes, which comes from the strong surface
band bending shown in the band diagram of Figure 5.6. As shown in Figure 5.10, the surface
electron density of the surface-charged device is considerably low at 1010 cm−2 orders compared
to 1013 cm−2 orders for the B-diffused device, resulting in high tolerance to the surface carrier
recombination in the surface-charged device structure. In other words, surface charged device
structure allows us to obtain high-efficiency solar cell performance even though there are large
amounts of interface states as high as ~1013 cm−2, which cannot be realized by B-diffused emitter
without reducing Nit as low as ~1010 cm−2 or further deposition of negatively charged dielectric
layers.

Fig. 5.9 Voc as a function of S for surface charge-induced emitter and B-diffused emitter.

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Fig. 5.10 Carrier densities for the devices with a surface charge-induced emitter or a Bdiffused emitter.
In addition, the free carrier absorption can be mitigated by the surface charge-induced emitter,obtaining slightly higher Jsc values. Free carrier absorption loss was calculated to be 0.035 mA
cm−2 for the B-diffused emitter and 0.003 mA cm−2 for the charged-induced emitter, respectively.
The free carrier absorption (αFC) was calculated using Green’s model [12], as expressed below.
𝛼FC = 2.6 × 10−27 𝑛𝜆3 + 2.7 × 10−24 𝑝𝜆2(5.2)

Where n is the electron density, p is the hole density, and λ is the wavelength. Parasitic absorption
can also be mitigated because SiO2 has no absorption (α = 0) over the wide wavelength of 210 to2000 nm [13], and the thickness is negligibly thin.
Furthermore, the surface charge-induced emitter shows large Auger lifetimes (τAuger) in the wide
region of the Si substrate due to shallow modulation doping. It suggests that Auger recombination
can be mitigated by a surface charge-induced emitter except for a very narrow region of less than
2 nm. The calculated depth profile of τAuger is shown in Figure 5.11. For the calculation, simulated
hole densities and the following equations [14] were used:
𝑅 = 𝐶 𝑛 𝑝2 ,

(5.3)

𝑛

(5.4)

1

𝜏Auger = 𝑅 = 𝐶𝑝2 .
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Where R is the recombination rate, and C is the Auger coefficient. For the value of C, 1.2×10−31
cm6 s−1 was used for p-type Si [15].

Fig. 5.11 Depth profile of τAuger for the devices with a surface charge-induced emitter or a Bdiffused emitter.
To summarize, the concept for the Al-COIL solar cell using surface charge-induced emitter canrealize high-throughput, low-cost p+ emitter formation by a simple wet chemical process,excellent ultrathin surface passivation with high tolerance to the interface states, reduction of free
carrier absorption loss, and no parasitic absorption using high band gap material of SiO2.

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5.3 Fabrication and evaluation of Al-COIL solar cells
5.3.1 Sample fabrication
The samples were fabricated on the double-side-polished FZ n-type Si (100) with a resistivity of
2.5 Ω cm and a thickness of 280 µm (TOPSIL PV-FZ™). The Si substrate was cleaved into 20×20mm-sized pieces, followed by RCA cleaning [16]. After 1% diluted hydrofluoric acid (HF)
treatment for 1 min and rinse in deionized water (DIW) (resistivity of 18.2 MΩ cm) for 1 min,the hydrophobic Si substrates were immersed in Al(NO3)3 aqueous solution with an
Al(NO3)3·9H2O molar concentration of 0.25 mol L−1 for 12.5 min. The hot Al(NO3)3 aqueous
solution was prepared in advance by dissolving Al(NO3)3·9H2O (17.8 g, special grade, Junsei
Chemical Co., Ltd.) in DIW (190 mL) in a Pyrex glass beaker (Corning Pyrex 600 mL) andheating to a boiling temperature (~95 °C) on a hot plate (AS ONE Corporation, CHPS-170DF).
During the heating process, the beaker was capped to avoid the evaporation of the solution. The
solutions were not intentionally stirred during the process; however, drastically generated bubbles
could be confirmed at the boiling temperature. After the Al(NO3)3 treatment, the processed Si
substrates were rinsed in DIW for 1 min and completely dried by N2 blow. Subsequently, the Ag
finger electrode was deposited on the Al(NO3)3-treated samples through a mask by thermal
evaporation in a vacuum chamber at a pressure of 9×10−4 Pa. For the rear side of the samples, thesurface was treated with a cotton swab dipped in 5% diluted HF aqueous solution in advance until
the hydrophobic surface was visually confirmed. The full-area Al electrodes were thermallyevaporated on the rear surfaces, which were scratched in advance to obtain an excellent ohmic
contact. The fabrication process flow and the fabricated solar cell structure are shown in Figure
5.12.

Fig. 5.12 Al-COIL solar cell fabrication process flow and the fabricated solar cell structure.
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5.3.2 Sample evaluation
The current density–voltage (J–V) curves were acquired in the dark and under 1 sun illumination(AM 1.5 G). The light intensity was calibrated with a Si reference cell (Bunkoukeiki Co., Ltd.,
BS-500BK Si photodiode detector) to an irradiance of 0.1 W cm−2 at RT. The Jsc was calculatedwith an active area of 3.67 cm2 after subtracting the estimated area of Ag front electrodes of 0.33
cm2. During the measurement, one probe was placed on the Ag finger electrode, and the other wasplaced on the Cu plate, which was contacted on the Al electrode of the rear side of the test cell tocollect the holes and electrons, respectively. The appearance of the setting and the photo J–V
measurement is shown in Figure 5.13.

Fig. 5.13 Appearance of the setting and the photo J–V measurement under 1 sun illumination.

For the discussion, photo J–V curves for the Qf dependency as front surface charges were
simulated using PC1D software, assuming the enormous front SRV of 1×1010 cm s−1 and without
front surface recombination. The other set of parameters is shown in the appendix. Additionally,τeff mapping measurement was performed after the one-side deposition of the Ag finger electrodeon the symmetrical structure of n-Si(100) with an Al-doped SiOx layer.
The τeff mapping data were acquired within the area of 20×20 mm2 with a 0.5 mm pitch. A pulsed
laser with a wavelength of 904 nm and an areal photon density of 5×1013 cm−2 was used for the
carrier excitation. The laser was illuminated on the rear side (without an Ag finger electrode) to
excite carriers evenly in the bulk of Si within the illuminated area.

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5.3.3

AL-INDUCED CHARGED OXIDE INVERSION LAYER SOLAR CELLS

J–V characteristics of Al-COIL solar cells

Figure 5.14 shows the J–V characteristics of a device with Al-doped SiOx, an Al-COIL solar cell,and a device without Al-doped SiOx acquired in the dark (dotted lines) and under 1 sunillumination (solid lines). As shown in Table 5.3 of the solar cell performance, a relatively high ηof 3.35% was acquired for the device with an Al-COIL structure. Although the η is quite low as ac-Si solar cell, a simple wet chemical process realizes the c-Si solar cells with a relatively highVoc of 0.454 V compared to the device without an Al-doped SiOx layer. Therefore, this result
demonstrated the device operation and the charge-induced inversion layer acts as a p+ emitter
formed by a simple wet chemical process. The low FF of 50.8% might be mainly attributed to theseries resistance, especially in a high sheet resistance of the front surface, but not from shunt
resistance because of no shunting occurrence in the dark J–V curve.

Fig. 5.14 J–V characteristics of the devices w/ Al-doped SiOx (Al-COIL solar cell) and w/oAl-doped SiOx acquired in the dark (dotted lines) and under 1 sun illumination (solid lines).

Table 5.3

Fabricated solar cell performance with or without Al-COIL structure.
Jsc (mA cm−2)

Voc (V)FFη (%)

w/ Al-doped SiOx (Al-COIL)

14.5

0.454

0.508

3.35

w/o Al-doped SiOx

2.9

0.140

0.528

0.21

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CHAPTER 5

AL-INDUCED CHARGED OXIDE INVERSION LAYER SOLAR CELLSAs described in Chapter 4, however, Al-doped SiOx/n-Si(100) showed a high implied Voc of
0.670 V, while the measured Voc of 0.454 V was relatively low. This might be mainly due to highrecombination at the rear surface receiving scratching and direct deposition of the Al electrode.
In addition, PC1D simulation shows that if the rear-SRV is as high as 1×1010 cm s−1, the simulated
Voc would be 0.587 V. Since this is still higher than the measured Voc of 0.454 V, further carrier
recombination might also occur at the front emitter surface. Figure 5.15 shows PC1D simulated
photo J–V characteristics for the devices with surface charge-induced emitters assuming various
Qf values in the range of −8.8×1010 – −1.0×1013 cm−2, enormous front-SRV of 1×1010 cm s−1, andemitter resistance of 5 Ω cm2 for each device. As shown in the figure, not only Voc but also Jsc are
significantly diminished with decreasing Qf values. Considering that the reduction of Jsc is notobserved without surface recombination, shown in Figure 5.16, the Jsc reduction might originate
from carrier recombination at the emitter surface due to increased minority carrier density in the
low Qf, as described in Figure 5.17.

Fig. 5.15 PC1D simulated photo J–V characteristics for the devices with surface chargeinduced emitters assuming various Qf values with a front S of 1×1010 cm s−1.

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AL-INDUCED CHARGED OXIDE INVERSION LAYER SOLAR CELLS

Fig. 5.16 PC1D simulated photo J–V characteristics for the devices with surface chargeinduced emitters assuming various Qf values without front-surface recombination.
.

Fig. 5.17 PC1D simulated carrier densities of devices with surface charge-induced emittersassuming low and high Qf values.

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The possible reason for the degradation is the reduction of the Vbi under the Ag finger electrodebecause Ag has a low work function: 4.26 eV for a polycrystalline phase and 4.52–4.74 eV for acrystalline phase [17], concerning the electron affinity of n-type c-Si (4.05 eV). Figure 5.18 showsτeff mapping measured after Ag finger deposition on the front side. The τeff mapping measurement
was performed from the rear side to avoid the suppression of the excess carrier generation andmicrowave reflection by the deposited Ag electrode. As seen in the τeff mapping, a significant
reduction of τeff has occurred at the deposited position. Hence, the deposited Ag electrode might
increase minority carrier (electron) density due to decreased band bending, which deteriorates thesurface passivation quality. Previous studies also showed that the metallization massively reducedτeff of ALD-Al2O3/SiO2/n-Si by Pd, Ni, and Al [18]. Although there is a contradiction to this work,
Ag did not show such degradation in the previous study [18]. However, higher work function
materials such as Pt, Ir, and Au should be applied to maintain the high Vbi. Favorably, relatively
low-cost transition metal oxides such as MoO3, WO3, and V2O5 (work function ~7 eV [19]) might
be used underneath the Ag electrodes.

Fig. 5.18 τeff mapping for Ag finger electrode deposited Al-doped SiOx/n-Si/ Al-doped SiOx/
structure and the schematic of the sample structure.

5.4 Summary
Al-COIL solar cells using surface charge-induced emitter realize high-throughput, low-cost p+
emitter formation by a simple wet chemical process. The PC1D simulation showed a highefficiency c-Si solar cell performance with a high tolerance to the interface states, a reduction offree carrier absorption, and Auger recombination in the wide region of the Si substrate. The first
fabricated test cell demonstrated a device operation with an η of 3.35% and a relatively high Voc
of 0.454 V, which implies that the induced inversion layer acts as a p+ emitter of an n-type c-Si
solar cell. However, a considerable divergence between Voc and iVoc was confirmed. In addition,
Jsc also diverged from the theoretical Jsc of ~26 mA cm−2. It might be mainly due to a combination
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of Qf degradation and high recombination at the rear and front surfaces. PC1D simulation revealed
that Voc and Jsc are significantly diminished with decreasing Qf values if the enormous front-SRV
is exceptionally high. The Qf reduction might have occurred under the Ag electrode due to fermilevel pinning. The promising solution is inserting a buffer layer between Ag and Al-doped SiOx
to mitigate the degradation; also, sheet resistance must be decreased by increasing Qf values in
the oxide layer.

References[1]

N. E. Grant, S. L. Pain, J. T. White, M. Walker, I. Prokes, and J. D. Murphy, “Enhanced
surface passivation of subnanometer silicon dioxide films by superacidic treatments,” ACS
Appl. Energy Mater. 5, 1542 (2022).

[2]

J.-I. Polzin, F. Feldmann, B. Steinhauser, M. Hermle, and S. Glunz, “Realization of
TOPCon using industrial scale PECVD equipment,” AIP Conf. Proc. 1999, 040018 (2018).

[3]

J. Haschke, O. Dupré, M. Boccard, and C. Ballif, “Silicon heterojunction solar cells: Recent
technological development and practical aspects - from lab to industry,” Sol. Energy Mater.
Sol. Cells 187, 140 (2018).

[4]

P. Van Halen, R. P. Mertens, R. J. Van Overstraeten, R. E. Thomas, and J. Van Meerbergen,“New TiOx-MIS and SiO2-MIS Silicon Solar Cells,” IEEE Trans. Electron Devices 25, 507(1978).

[5]

R. E. Thomas, R. B. North, and C. E. Norman, “Low cost - high efficiency MIS/inversion
layer solar cells,” IEEE Electron Device Lett. 1, 79 (1980).

[6]

F. Werner, Y. Larionova, D. Zielke, T. Ohrdes, and J. Schmidt, “Aluminum-oxide-based
inversion layer solar cells on n-type crystalline silicon: Fundamental properties andefficiency potential,” J. Appl. Phys. 115, 073702 (2014).

[7]

K. Jager, and R. Hezel, “Optical stability of silicon nitride MIS inversion layer solar
cells,” IEEE Trans. Electron Devices 32, 1824 (1985).

[8]

M. Yu, Y. Shi, J. Deru, I. Al-Dhahir, S. McNab, D. Chen, M. Voss, E.-T. Hwu, A. Ciesla,
B. Hallam, P. Hamer, P. P. Altermatt, P. Wilshaw, and R. S. Bonilla, “Assessing the potential
of inversion layer solar cells based on highly charged dielectric nanolayers,” Phys. Status
Solidi RRL 15, 2100129 (2021).

[9]International Technology Roadmap for Photovoltaic (ITRPV) 2022 Results 14th edition(2023).

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J. Bao, C. Chen, L. Ma, C. Huang, X. Zhang, Z. Wang, C. Chen, S. Zhan, R. Liu, Z. Qiao,
Z. Du, Z. Liu, and J. Chen, “Towards 24% efficiency for industrial n-type bifacial
passivating-contact solar cells with homogeneous emitter,” In: Proceedings of the 37th
European PV Solar Energy Conference and Exhibition, Online, 160–163 (2020).

[11]

J. Schmidt, R. Peibst, and R. Brendel, “Surface passivation of crystalline silicon solar cells:
Present and future,” Sol. Energy Mater. Sol. Cells 187, 39 (2018).

[12]

M. A. Green, Silicon Solar Cells: Advanced Principles & Practice (University of New
South Wales, Sydney, Australia 1995).

[13]

E. D. Palik, Handbook of Optical Constants of Solids Ⅰ (Academic Press, 1985).

[14]

A. Goetzberger, J. Knobloch, and B. Voß, Crystalline Silicon Solar Cells (John Wiley &Sons, 1998).

[15]

J. D. Beck and R. Conradt, “Auger recombination in Si,” Solid State Commun. 13, 93(1973).

[16]

W. Kern and D. A. Puotinen, “Cleaning solution based on hydrogen peroxide for use insilicon semiconductor technology,” RCA Rev. 31, 187 (1970).

[17]

A. W. Dweydari and C. H. B. Mee, “Work function measurements on (100) and (110)
surfaces of silver,” Phys status solidi (a) 27, 223 (1975).

[18]

D. Hiller, P. Hönicke, and D. König, “Material combination of tunnel-SiO2 with a(sub-)monolayer of ALD-AlOx on silicon offering a highly passivating hole selective
contact,” Sol. Energy Mater. Sol. Cells 215, 110654 (2020).

[19]

M. T. Greiner, L. Chai, M. G. Helander, W.-M. Tang, and Z.-H. Lu, “Transition metal oxide
work functions: The influence of cation oxidation state and oxygen vacancies,” Adv. Funct.
Mater. 22, 4557 (2012).

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ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

CHAPTER 6
ANTI-REFLECTIVE STRUCTURES AND
LAYERS ON AL-COIL ARCHITECTURE

In the first test cell, it was discovered that the surface charge-induced inversion layer acted as ap+ emitter for an n-type c-Si solar cell. However, the solar cell performance was significantly
lower than that of commercial c-Si solar cells. The front surface of the conventional c-Si solar
cell has pyramidal textured Si structures to increase light-trapping, boosting photocurrent density.
Moreover, to increase photocurrent density further, an ARC, such as the SiNy layer with anadequate refractive index, is deposited on the pyramidal textured Si surface. According to theoptical simulation, a photocurrent density of 14.75 mA cm−2 will be increased by applying
optically optimized SiNy and pyramidal textured Si substrate. Since this increment is almost
double for the first test cell, implementing such anti-reflective structures and layers will be a
promising solution to improve the performance of Al-COIL solar cells. In addition, an increased
carrier generation rate might enhance Qf in the Al-doped SiOx layer for Al-COIL solar cells,
leading to further hole-selective passivating contact properties. Hence, anti-reflective structuresand layers possibly also increase Voc and FF.
In this chapter, the effectiveness of the anti-reflective structures and layers on Al-COIL
architecture is elucidated to access high-efficiency c-Si solar cells, from the fabrication ofpyramidal textured Si surfaces to the evaluation of the solar cell performance, including the studyof SiNy/Al-doped SiOx stacks for testing the surface passivation quality and optical property.

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6.1 Strategy of improvement for Al-COIL solar cell performance
6.1.1

Enhancement of photocurrent density

A pyramidal textured structure is conventionally formed on the c-Si surface to enhancephotocurrent density due to increased light-trapping in the c-Si substrate. Since the reflected lightfrom the textured Si surface is reduced by increasing the chances of reflected light bouncing returnto the surface, rather than out to the air, increasing light trapping. Table 6.1 shows the photon
current of the incident (Jinc), reflected (JR), absorbed in film (JA), and absorbed in the substrate(JG) under 1 sun illumination, simulated by OPAL 2 simulator [1]. PECVD-SiNy (n = 1.99) wasused as an ARC, and the Si substrate thickness was set to 200 µm. Lambertian limit [2] wasconsidered to be a light-trapping model. The simulation shows optimized ARC on the pyramidal
textured Si substrate, which realizes JG = 43.64 mA cm−2. In contrast, planar Si substrate without
ARC shows 28.89 mA cm−2. Therefore, the combination of ARC and pyramidal texture increases
carrier generation significantly, and implementing this on the Al-COIL structure can lead to a 1.51
times enhancement of photocurrent density. The anti-reflective effect of pyramidal texture and
ARC on the Si substrate is depicted in Figure 6.1.

(a)Fig. 6.1(b)Anti-reflective effect by (a) pyramidal textured Si surface and (b) anti-reflective

coating (ARC) on Si substrate.
Table 6.1

Photon current density (JG) simulated by OPAL 2.
Jinc (mA cm−2)

JR (mA cm−2)

JA (mA cm−2)

JG (mA cm−2)

w/o ARC planar sub.

44.54 (100%)

15.65 (35.1%)

0.00 (0.0%)

28.89 (64.9%)

w/o ARC texture sub.

44.54 (100%)

4.28 (9.6%)

0.04 (0.1%)

40.22 (90.3%)

w/ ARC planar sub.

44.54 (100%)

4.47 (10.0%)

0.00 (0.0%)

40.07 (90.0%)

w/ ARC texture sub.

44.54 (100%)

0.83 (1.9%)

0.07 (0.2%)

43.64 (98.0%)

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6.1.2

ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

Enhancement of surface passivation quality

As described earlier, the enhancement of carrier generation up to 1.51 times will be theoreticallyrealized compared to the planar Si surface by implementing the pyramidal textured surface with
ARC. As expressed in the equation below [3], in the sufficiently high negative charge condition,
the Seff will be decreased with increasing surface hole density (ps). Since ps is proportional tosquared Qf, the Seff will be dramatically decreased by the Qf enhancement.
𝑆𝑛0 𝑛d
𝑝s

(6.1)

𝑄f 2
𝑝s =
2𝑘B 𝑇𝜀Si(6.2)

𝑆eff =

Where Sn0 is the SRV of the electron, and nd is the sum of Ndop and Δn.
At the same time, the increased excited electrons expectedly trapped in the Al-induced acceptor
states of the Al-doped SiOx layer, which leads to a higher negative Qf value with increasing excess
carriers at the surface that might be in the inversion layer and the generated carrier at the outermost
surface might dominantly affect the carrier trapping. Thus, the depth profile of the generation rate(G) was simulated by the OPAL 2 simulator. As shown in Figure 6.2 (a), the pyramidal texture
and pyramidal texture with ARC show higher G values of 2.31×1014 and 3.05×1014 cm−2 s−1,
respectively. Figure 6.2 (b) shows the ratio of the G values concerning the G value of the planar
Si substrate. 1.4 times higher generation occurs at the outermost surface of pyramidal textured Si
substrate even without ARC. Hence, a pyramidal textured surface is expected to induce more Qfvalue. Assuming 1.4 times increment of Δn, the Qf value enhances 1.4 times if the excess carriers
are entirely trapped in the Al-induced acceptor states. Additionally, the steeper angle of the highaspect ratio of pyramidal texture with a steeper angle of 80° shows higher G values. It is 1.65
times higher than that of the planar surface. This significant enhancement of G values is attributed
to the excellent light-trapping by a pyramidal texture with a pillar-like structure with a high aspectratio, which can realize ultra-low reflection of only 0.2% in the whole wavelength range of 300
to 1200 nm, as shown in Figure 6.3. Therefore, in this type of solar cell, light absorption issignificant for not only Jsc but also Voc.

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ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE(b)Fig. 6.2 (a)Depth profile of the generation rate (G) for the Si substrates with planar or

pyramidal texture surface w/ or w/o ARC, including a pyramidal texture with a base angle of 80°.
(b) The ratio of the G values concerning the G value of planar Si substrate.

Fig. 6.3

Reflectance spectra for the Si substrates with planar or pyramidal texture surface w/ or

w/o ARC, including a pyramidal texture with a base angle of 80°.

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As in the following equations [3], Seff will be squared decreased with increasing Q value. Hence,in the case of ARC-coated pyramidal textured Si substrate, the Seff and the surface saturation
current density (J0S) will be reduced to 1.96 times lower than the planar Si surface.
2𝑘B 𝑇𝜀Si 𝑆n0 𝑛d
𝑄2

(6.3)

2𝑘B 𝑇𝜀Si 𝑆n0
𝑄2

(6.4)

𝑆n0 = 𝜈thn 𝑁it 𝜎n(6.5)

𝑆eff =

𝐽0S = 𝑞 𝑛i2

Where T is the absolute temperature, νthn is the thermal velocity of the electron, Nit is the interface
state density per area, and σn is the capture cross-section of the electron, respectively.
However, the pyramidal textured Si surface has a 1.73 times larger surface area than the planar
surface. Therefore, the increment of the surface area should be considered to influence the Nit
value. If the Nit value is proportionally increased with increasing the surface area, Seff and J0S will
be reduced ~1.13 times lower than that of the planar Si substrate.
In addition, as shown in Figure 6.2, since Rsheet reduces inversely proportional to the Qf2, Rsheet
possibly reduces ~50% by 1.4 times increment of the carrier generation. Therefore, Jsc, Voc, andFF are expected to be improved using pyramidal texture with ARC by this light-trapping approach.
However, the front emitter surface should be sufficiently low Rsheet, as low as less than 100 Ω sq−1.
To enable such a low level of Rsheet, a negative Qf of ~6×1013 cm−2 might will be needed,
corresponding to a very high ps of ~1×1022 cm−3.

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6.2

ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

Effectiveness of pyramidal textured Si surfaces

The pyramidal textured Si surface typically has a sub-micrometer-sized pyramidal structure witha base angle of 54.74°, as shown in Figure 6.4. The structure consists of several crystal
orientations: (100) for the top and bottom, (111) for the slope, and (110) for the ridgeline. Since
most of the surface contains (111)-facet, c-Si solar cell performance should have a high-quality
hole-selective passivating contact on a (111) surface. In general, however, the pyramidal textured
Si makes a tiny structure at the top and the bottom of the valley, which causes insufficient surface
cleaning due to increased surface tension of the cleaning solution [4] and degradation of thefollowing surface passivation. Hence, to improve the passivation quality on pyramidal textured
Si, it would be crucial to reduce the surface tension.

Fig. 6.4 Structure of pyramidal textured Si surface.
However, the field-effect passivation effectively mitigates the surface recombination by thehighly negative charges despite the large surface area on nano-sized structures due to the increaseof hole density and the reduction of electron density. The previous study showed that black-Si (bSi) surface can be effectively passivated by a highly negatively charged ALD-Al2O3 anddemonstrated more than 22% high-efficiency b-Si solar cells [5, 6]. They mentioned the excellent
passivation mechanism is attributed to the enhancement of the effective Qf values by the larger
effective surface area. Since the surface charges affect the sub-hundred-nanometer region to
generate mirror charges, the opposite face surface charge is affected due to the distance of the
sub-hundred nanometer between both sides. The dependence of the distance from the opposite
side was simulated using PC1D simulation to investigate the effect of the opposite surface of thenegative charges. Figure 6.5 shows the depth profile of carrier densities with highly negative Qf
of 1×1013 cm−2 for both surfaces. As shown in the Figure, hole density gradually increases, but
electron density significantly decreases with closing the faced-surface distance. In contrast, thedistance sufficiently far away from each surface does not affect carrier density in the Si substrate.
The high Qf in the oxide generates positive charge carriers in the space charge region as mirroredcharges. If the space charge region was compressed by reducing the face-to-face distance of Si
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surfaces, the space charges are increased when the surface Qf is moderate. Hence, the simulation
implies that a nano-sized structure with sufficiently high negative Qf enhances hole density andreduces minority carriers in the Si surface. This enhancement originated from the space charge
region compression effect [7,8]. The schematics of the face-to-face structures with nano-sized or
micrometer-sized distances between both sides with negative charges are shown in Figure 6.6. In
the case of a nano-sized structure, induced mirror charges exist in bulk Si. On the other hand, inthe case of a micrometer-sized structure, the mirrored charges only exist at the surface of Si. The
nano-sized structure induces an inversion layer in the whole Si structure, as shown in Figures 6.5
and 6.6. Hence, the nano-sized pyramidal textured Si surface with highly negative charges
expectedly enhances the hole-selective passivating contact. However, the typical alkaline
texturing process is unavailable for fabricating such a nano-sized structure. As the first stage, I
focused on the effectiveness of conventionally available pyramidal texture formed by alkaline
texturization for the Al-COIL solar cells.
(a)

(b)

(c)

(d)

Fig. 6.5 Depth profile of carrier densities of face-to-face structure with distances of (a) 10 nm,(b) 50 nm, (c) 100 nm, and (d) 1 µm between both sides with highly negative Qf of 1×1013 cm−2.

Fig. 6.6

Schematics of the face-to-face structures with nano-sized or micrometer-sized distance

between both sides with negative charges.
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6.3 Sample fabrication
6.3.1

Texturing process

Double-side-polished 280 µm thick FZ n-type Si (100) with a resistivity of 2.5 Ω cm (TOPSIL
PV-FZ™) was used. The Si substrate was cleaved into 50 × 50 mm-sized pieces and cleaned byultrasonication with acetone (Electrical grade, Kanto Chemical Co., Inc.) followed byultrasonication with ethanol (Electrical grade, Kanto Chemical Co., Inc.). The cleaned samples
were immersed into an alkaline solution at 80 °C for 15 min with a magnetic stirrer agitating (600rpm) to form pyramidal textures on Si substrates. The 80 °C alkaline solution was prepared inadvance by dissolving 85% potassium hydroxide (KOH) (9 g, Nacalai Tesque Inc.) into DIW (440mL) with an addition of TP101 (1.35 g, Hayashi pure chemical Ind., Ltd.) as a micro-masking
agent.

6.3.2

Cleaning and surface passivation

After the texturing process, the samples were rinsed with DIW for >2 min. After the treatment of
5% diluted hydrofluoric acid (DHF) containing 60% methanol for 2 min and rinse in DIW for 1
min, the textured Si substrates were cleaned by RCA cleaning. Between SC-1 and SC-2, and after
the SC-2 cleaning, the methanol-added DHF treatment was performed for 2 min. After rinsed with
DIW for 1 min followed by ethanol for 1 min, Al(NO3)3 treatment was performed in Al(NO3)3
aqueous solution with Al(NO3)3·9H2O molar concentration of 0.25 mol L−1 for 12.5 min. The hot
Al(NO3)3 aqueous solution was prepared in advance by dissolving Al(NO3)3·9H2O (28.1 g, special
grade, Junsei Chemical Co., Ltd.) and added polyethylene glycol (PEG) (3 mL, molecular weightof 200, extra pure reagent, Junsei Chemical Co., Ltd.) as a surfactant in DIW (300 mL) in a Pyrex
glass beaker (SCI 1000 mL) and heated to a boiling temperature (~95 °C) on a hot plate (AS ONE
Corporation, CHPS-170DF). During the heating process, the beaker was capped to avoid theevaporation of the solution. The solutions were not intentionally stirred during the process;however, drastically generated bubbles could be confirmed at the boiling temperature. After the
Al(NO3)3 treatment, processed Si substrates were rinsed in DIW for 1 min, followed by ethanolfor 1 min, and entirely dried by N2 blow.
The sample fabrication process flow, the sample structure, and the appearance of the Al(NO3)3
treatment are shown in Figures 6.7 and 6.8, respectively.

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Fig. 6.7 Sample fabrication process flow

(b)

(a)

Fig. 6.8 (a) Sample structure and (b) appearance of Al(NO3)3 treatment.

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6.4 Sample evaluation
The τeff measurements were performed using μ-PCD (Kobelco Research Institute, Inc., LTA1510EP). The τeff mapping data were acquired within the 44×44 mm2 area in 50×50 mm2 sized
samples with a 1 mm pitch. A pulsed laser with a wavelength of 904 nm and an areal photon
density of 5 × 1013 cm−2 was used for the carrier excitation. Additionally, τeff and iVoc as a functionof Δn were analyzed using QSSPC (Sinton Instruments, WCT-120). For the analysis, theestimated value of 1.05 was used as an optical factor (fopt). Using the following relationship, thefopt was calculated from photon current density (JG) of 40.03 mA cm−2 generated from measured
reflectance and transmittance spectra. The reflectance and transmittance spectra were acquired
using a UV-vis-NIR spectrophotometer (Shimadzu, UV-3150).

𝐴(𝜆) = 1 − 𝑅(𝜆) − 𝑇(𝜆)

𝜆

(6.6)

𝐽G = 𝑞 ∫𝜆 1200 𝐹ph (𝜆) 𝐴(𝜆) d𝜆(6.7)

𝐽 (mA cm−2 )
38 (mA cm−2 )

(6.8)

300

𝑓opt = G

Here, A(λ), R(λ), T(λ), and Fph(λ) are absorbance, reflectance, transmittance, and photon flux ateach wavelength. JG is photon current density, and q is elemental charge. 38 mA cm−2 means thephotocurrent density of a standard cell, which is mounted in the Sinton lifetime tester, and the
38 mA cm−2 corresponds to fopt = 1. The Seff, max was calculated using Eq. (4.3). Since the Seff, max
calculation requires wafer thickness, it is measured using a micrometer at the 9 points of samples
after τeff measurements. The measurements estimated 269 µm as an average thickness. The
measured thicknesses are shown in Figure 5.9. In addition, the J0 value was also analyzed by theKane and Swanson method using the auger-corrected effective minority carrier lifetime (τcorr)
over the Δn range of 1×1015 to 9×1015 cm−3.

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Fig. 6.9 Measured wafer thickness of a fabricated pyramidal textured Si surface.

The surface morphology was observed by SEM (Hitachi, TM3030Plus) and AFM mode,
mounted in SPM (Bruker, Dimension Icon). Additionally, the surface potential was measured
by a kelvin probe force microscope (KPFM) mounted in the SPM system. Pt-Ir coated Sb highly
doped Si (Bruker, SCM-PIT) was used for the tip of the probe. The sample was placed on thesample stage with an Ag paste on the corner of the sample to form an electrical contact between
the sample and the sample stage.

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6.5 Surface morphology
Figure 6.10 shows the SEM image of the fabricated pyramidal textured Si surface. The ~2 µm
sized pyramidal textures were randomly formed on the Si surface, and less than 1 µm sizedpyramidal texture also existed on the Si substrate. The planar surface was transformed entirely
into pyramidal structures, and the shape of the tip, valley, edge, and face structures was seen. Most
structures consist of (111)-facets at the face structures. AFM measurement revealed the depth ofpyramidal texture is ~1 µm, as shown in Figure 6.11. The estimated base angle was a lower valueof 49.8°compared with a theoretical value of 54.74°.

Fig. 6.10 SEM image of the fabricated pyramidal textured Si surface.

Fig. 6.11 AFM image and the depth profile of the fabricated pyramidal textured Si surface.

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ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

Estimation of photon current density

Figure 6.12 shows (a) the reflectance (R), (b) transmittance (T) spectra, and Figure 6.13 shows
the calculated absorbance (A) spectra for the fabricated pyramidal textured Si substrate and mirror
polished Si substrate as a reference. The pyramidal textured Si surface significantly decreased
reflectance below 10% in the wide wavelength range of 500–1070 nm. In addition, thetransmittance also reduced in the wavelength range of ~950–1200 nm. The calculated absorbance
spectra estimated the JG to be 40.03 mA cm−2 for pyramidal texture and 26.00 mA cm−2 for planarusing eq. (5.6)–(5.8), respectively. Hence, the JG gain from the pyramidal textured Si substrate
was estimated to be 14.03 mA cm−2. The calculated JG and fopt are summarized in Table 6.2.

(a)

(b)

Fig. 6.12 (a) Reflectance and (b) transmittance spectra for the fabricated pyramidal textured
and mirror polished Si substrate.

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Fig. 6.13 Calculated absorbance spectra for the fabricated pyramidal textured and mirrorpolished Si substrate.

Table 6.2

Estimated photon current density (JG) and optical factor (fopt).
Substrate

JG (mA cm−2)

fopt

Planar

26.00

0.69

Pyramidal texture

40.03

1.05

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6.7 Surface passivation quality
Figure 6.14 shows time-related τeff mappings for the passivated pyramidal textured Si substrate
stored in ambient air. The high τeff could not be obtained after the passivation and decreased after
10 min. However, after 6 h stored in the ambient air, the τeff was dramatically increased to 491 µs,
on average. This time-related τeff increase might be attributed to the lack of surface oxidation. If
the formed interfacial SiOx is too thin, the Qf value will be decreased due to the low concentrationof AlO4 structure. Thus, the interfacial Si will oxidize while stored in ambient air, resulting inhigher surface passivation quality with a negative Qf value. For this sample fabrication, methanol
rinse was not performed before the Al(NO3)3 treatment, and PEG was not added to the Al(NO3)3
aqueous solution. The pyramidal texture has very fine structures, especially in the part of thevalley. Therefore, it might be challenging to passivate such a fine structure due to low wettabilityon the surface [9].

Fig. 6.14 Time-related τeff mappings acquired by µ-PCD for the passivated pyramidal textured
Si substrate stored in ambient air.

To enhance the wettability and oxidation reaction on the pyramidal textured structure, one ofthe promising solutions was performing an ethanol rinse before the Al(NO3)3 treatment to remove
the residue of HF chemicals altogether. Moreover, adding a surfactant into Al(NO3)3 aqueous
solution was another solution. Therefore, the ethanol rinse was performed before the passivation
and Al(NO3)3 treatment with PEG as a surfactant. Figure 6.15 shows τeff mappings acquired by µPCD for (a) just after the surface passivation and (b) after 6 hours. As a result, good surface
passivation quality was observed just after the passivation. On average, τeff of 400 µs and 474 µs,
and 698 µs and 692 µs were obtained as maximum τeff, respectively. As described in Chapter 4,
the Al(NO3)3-treated mirror polished n-type Si(111) showed a τeff of 360 µs. If compared just after131

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the passivated sample, passivated pyramidal textured Si shows slightly higher τeff. However, timerelated changes were confirmed in the τeff. After 6 h storage in ambient air, the low τeff region
became higher τeff. On the other hand, the high τeff region became slightly lower τeff.

(a)

(b)

Fig. 6.15 τeff mappings acquired by µ-PCD for (a) just after the surface passivation and (b) after
6 hours of storage in ambient air.

Figure 6.16 shows τeff and Seff, max of the pyramidal textured Si substrate treated in boiled 0.25
mol L−1 Al(NO3)3 (aq.) for 12.5 min as a function of Δn for just after the surface passivation andafter 6 hours storage in ambient air. Considering the low Δn region of ~1014 cm−3 shows highsurface passivation quality as τeff of >1 ms, the field effect might have dominantly originated fromsurface passivation. The high τeff of 935 and 1012 µs were confirmed at Δn =1×1015 cm−3, and thecorresponding Seff, max were estimated to be 14.4 and 13.3 cm s−1, respectively. The iVoc were
estimated to be 680 and 682 mV at 1 sun illumination, and J0 were analyzed as 68 and 59 fA cm−2,
respectively. The fitted Auger-corrected inverse effective lifetimes and iVoc as a function of lightintensity are shown in Figures 6.17 and 6.18, respectively.

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τeff (ms)CHAPTER 6

Fig. 6.16

τeff and Seff, max of the pyramidal textured Si substrate treated in boiled 0.25 mol L−1

Al(NO3)3 (aq.) for 12.5 min as a function of Δn for just after the surface passivation and after 6

1/τeff −1/τAuger (s−1)

hours storage in ambient air.

Fig. 6.17

Fitted Auger-corrected inverse effective lifetimes of the pyramidal textured Si

substrate treated in boiled 0.25 mol L−1 Al(NO3)3 (aq.) for 12.5 min as a function of Δn for justafter the surface passivation and after 6 hours storage in ambient air.
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Fig. 6.18 iVoc of the pyramidal textured Si substrate treated in boiled 0.25 mol L−1 Al(NO3)3(aq.) for 12.5 min as a function of light intensity for just after the surface passivation and after 6hours storage in ambient air.

For pyramidal textured Si substrate, it is difficult to evenly passivate Si surface with a large
area of 50×50 mm2 at a laboratory level. It might be due to unevenness of the process temperatureduring the Al(NO3)3 treatment because the heating process was performed using a hot plate, which
can only heat from the bottom of a beaker. Therefore, 15×15 mm2 sized relatively small Si pieces
are evaluated to compare the surface passivation quality of the planar and pyramidal textured Si
substrates. Figure 6.19 shows the τeff mappings acquired by µ-PCD for an Al(NO3)3-treatedpyramidal textured Si substrate and a planar Si (111) substrate. These pieces are proceeded
through the same process (cleaning and surface passivation). The pyramidal texture showed asignificant increase in τeff up to 688 µs on average compared with a planar substrate of 387 µs.
The 1.77 times higher τeff might have originated from the increased carrier generation. The
estimated absorbances of the 904 nm light are 93.4% for a textured surface and 68.1% for a planar
surface, respectively. Hence, the pyramidal textured Si substrate can absorb a 1.37 times higher
photon concentration. If the increased excess carriers can be trapped in the Al-induced acceptor
states, τeff will be increased 1.88 times because τeff is proportional to Qf2 in sufficiently highnegative charged conditions. Since this enlargement is close to the τeff increment, the τeff increment
mainly originated from an increase in light-trapping. However, considering Dit increase by thesurface area of pyramidal texture, another effect must be regarded to explain this τeff increment.
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(a)

(b)

Fig. 6.19

τeff mappings acquired by µ-PCD for the Al(NO3)3-treated (a) pyramidal textured and(b) planar Si substrate.

6.8 Surface potential
The surface potential of Al-COIL on a pyramidal textured Si surface was investigated using
KPFM, which is mounted in the SPM system. Figures 6.20 (a), (b), and (c) show the heightmapping, and Figures 6.20 (d), (e), and (f) show the surface potential of the planar, valley regionof pyramidal texture, and tip region of pyramidal texture, respectively. The planar surface andvalley region of pyramidal texture showed a comparable surface potential of −0.72 V and −0.75
V, respectively. However, the tip region of pyramidal texture showed a lower surface potential of−0.98 V. Since the used probe was Pt-Ir coated Sb highly doped Si with a work function of 4.86
V [9], which is very close to the valence band edge of Si, Al-doped SiOx formed n-type Si surface
bands are strongly bent upward, especially in the tip region of the pyramidal textured surface. The
tip of a pyramidal textured Si surface can be recognized as a nano-sized structure due to opposite
side surfaces with highly negative charges within a nano-scale distance. Therefore, the lower
surface potential on the pyramid's tip demonstrated that such a nano-sized structure enhances holeconcentration, causing strong upward band bending.

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(b)

(a)

30 nm

(c)

−30 nm1 µm(d)

100 nm

350 nm(e)−0.6 V(f)−1.2 V
1 µm

350 nm

100 nmFig. 6.20 (a), (b), and (c): height mapping and (d), (e), and (f): surface potential for the planar,
valley region of pyramidal texture, and tip region of pyramidal texture, respectively.

Figure 6.21 shows simulated 2D mapping of surface hole density on a pyramidal textured Si
surface with a high negative Qf of 3.0×1013 cm−2 using PC3S [10] simulation. The simulation was
performed under 1 sun illumination for a 2 µm size pyramidal texture with a base angle of 54.74°
on a 15 µm thick n-type Si substrate with a doping concentration of 2.0×1015 cm−3. This 2D
simulation suggests that surface hole density on the tip and edge of the pyramidal structure isenhanced. Thus, 2D simulation assists the lower surface potential of the tip of the pyramidal
structure. Therefore, sub-micrometer-sized typical pyramidal textured Si substrates have opposite
side surfaces with highly negative charges within a nano-scale distance, which enhances thesurface hole density. In such sufficiently high negative Qf conditions, Seff will be inversely
proportional to the surface hole density, which realizes high surface passivation quality and holeselectivity. However, in such high negative Qf conditions, the averaged surface hole density isalmost the same as a planar surface. This would be one of the reasons for the higher surface
passivation quality on a pyramidal textured Si surface in the low negative Qf conditions, such as
Qf = 1011cm−2, a significant increase can be seen from 9.99×1010 to 4.80×1014 cm−3 by applying
pyramidal texture, as shown in Figure. 6.22. When the estimated Qf value was considered as−3.2×1012 cm−2, pyramidal texture showed a slight increase of surface hole density from 9.52×1019
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to 1.08×1020 cm−3. Assuming most of the surface orientations are (111), Dit increase is attributed
to the ~1.73 times larger surface area of the pyramidal texture. Therefore, the 1.13 times higher
surface hole density might not compensate for the increased Dit. Hence, the increase of carrier
generation rate by enhanced light-trapping dominates the improvement of the surface passivation
quality.

1.84×1021

5.42×1021(cm−3)
Fig. 6.21

PC3S simulated 2D mappings of surface hole density on (a) pyramidal textured and(b) planar Si surface with a high negative Qf of 3.0×1013cm−3. The color scale is shown inpercentile; the center value is shown as yellow.

2.61×109

2.97×1016(cm−3)
Fig. 6.22

PC3S simulated 2D mappings of surface hole density on (a) pyramidal textured and(b) planar Si surface with a high negative Qf of 1.0×1011 cm−3. The color scale is shown inpercentile; the center value is shown as yellow.
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6.9 Fabrication and evaluation of solar cells
6.9.1 Sample fabrication
The Ag finger electrode was deposited on the Al(NO3)3-treated sample through a mask by thermal
evaporation in a vacuum chamber at a pressure of 9×10−4 Pa to evaluate solar cell performance.
To form a grid electrode, after the deposition of the Ag finger electrode, the deposited sample wasrotated 90°; subsequently, the Ag finger electrode was deposited. For the rear side of the sample,
the full-area Mg followed by Al electrodes were thermally evaporated on the rear surfaces. The
rear side of the Si band bends downwards because Mg is a low work function (3.66 eV [11])material concerning n-type Si, resulting in the formation of electron selective contact by anaccumulation layer. The fabrication process flow and the fabricated solar cell structure are shownin Figures 6.23 and 6.24, respectively.

Fig. 6.23 Fabrication process flow of Al-COIL solar cell with a pyramidal textured surface.
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Fig. 6.24 Fabricated Al-COIL solar cell structure and front Ag grid electrode.

6.9.2 Sample evaluation
The current density–voltage (J–V) curves were acquired in the dark and under 1 sun illumination(AM 1.5 G). The light intensity was calibrated with a Si reference cell (Bunkoukeiki Co., Ltd.,
BS-500BK Si photodiode detector) to an irradiance of 0.1 W cm−2 at RT. The Jsc was calculatedwith an active area of 0.251 cm2 after subtracting the estimated area of Ag front electrodes in an
opened area of black-shade mask of 0.624 cm2. During the measurement, one probe was placedon the Ag finger electrode, and the other was placed on the Cu plate, which was contacted on the
Al electrode of the rear side of the test cell to collect the holes and electrons, respectively.

6.9.3

J–V characteristics of Al-COIL solar cells

Figure 6.25 shows J–V curves of the devices with AL-COIL structures on a pyramidal texture anda planar surface acquired in the dark (dotted line) and under 1 sun illumination (solid line). The
solar cell performances are shown in Table 6.3. As a result, the device fabricated on a planar Si
substrate led to a low η of 2.80%. In comparison, a significantly improved η of 12.13% was
acquired for the device fabricated on a pyramidal textured Si substrate. In particular, a very highFF of 0.794 was achieved. The high FF was realized by a fine grid Ag electrode on an emitter sidewith a larger surface area of pyramidal textured surface. This can mitigate a low lateral
conductivity in the emitter induced by Al-COIL. However, the Voc is still lower than the analyzed
iVoc of 682 mV. The low Voc might be attributed to the Vbi reduction under the Ag grid electrodedue to a low work function of Ag, as discussed in Chapter 4. In addition, the Jsc also showed asignificant increase, but still lower than the theoretical value of 40.03 mA cm −2. This Jsc139

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degradation might have originated from the carrier recombination at the front emitter surface, as
discussed in Chapter 4. If there are many carrier recombination centers at the emitter side and adegradation of charged oxide as a low negative Qf, the Jsc degradation occurs as shown in the J–
V curves of Figure 6.26 acquired from PC1D simulation.

Fig. 6.25 J–V curves of the fabricated solar cells with Al-COIL structures on a pyramidal texture
and a planar surface.

Table 6.3

Solar cell performances of the AL-COIL on a pyramidal texture and a planar surface.
Jsc (mA cm−2)

Voc (V)FFη (%)

Pyramidal texture

29.63

0.516

0.794

12.13

Planar n-Si(111)

21.50

0.210

0.622

2.80

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Fig. 6.26

ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

PC1D simulated Jsc as a function of surface recombination velocity (SRV) withvarying surface Qf values.

6.10

Anti-reflective coating on Al-COIL structure

One possible solution to increase the Jsc of Al-COIL solar cells is through the use of ARC, which
can enhance light trapping by gradually changing the refractive index (n) between air (n=1) andthe Si substrate (n = 3.88 at a wavelength of 630 nm) [12]. The reflection can be calculated using
the Fresnel equations. If the light incidents are perpendicular to the surface, the reflectance isexpressed by𝑛 −𝑛

𝑅 = (𝑛1 +𝑛2 )2
1

2

(6.9)

Where n1 is the refractive index of the first media and n2 is the refractive index of the second
media, respectively.
The optimal value of n is produced by the geometric mean of the two surrounding indices, as
described below.
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𝑛ARC = √𝑛air 𝑛sub.

(6.10)

Therefore, by inserting a layer with an n value of 1.97, the reflection of light can be mitigated,
and SiNy is commonly used for this purpose due to its n value. Furthermore, the key to reducing
reflection is to ensure that the reflected waves from these two interfaces between air/ARC and
ARC/Si have a path difference of λ/2 (half-wavelength). This path difference is crucial for
destructive interference to occur, where the peaks of one wave align with the troughs of the other,resulting in minimal reflected light. To achieve a path difference of λ/2, the film should be λ/4
thick, corresponding to a quarter-wavelength. Hence, the optimal thickness of ARC is given by
the equation below.

𝜆𝑑ARC = 4𝑛

ARC(6.11)

Hence, the thickness of ARC should be 80 nm to minimize the reflection of light. Typically, SiNy
is deposited using plasma-enhanced chemical vapor deposition (PECVD), but this method candamage ultrathin layers on Si substrates and affect surface passivation quality and solar cell
performance [13]. As an alternative, catalytic chemical vapor deposition (Cat-CVD) offers agentler approach to deposit, using catalytically generated radicals on a hot wire with a hightemperature of approximately 1800 °C, without plasma. A previous study demonstrated the CatCVD-grown SiNy layer has an optical transparency with n~2 and excellent surface passivation
quality [14]. Therefore, SiNy /Al-doped SiOx stacks were investigated to determine whether the
stacks show excellent surface passivation quality and optical properties.

6.10.1 Sample fabrication
Double-side-polished n-type FZ Si(100) wafers with a thickness of 280 μm, a resistivity of 2.5
Ωcm, and a size of 50×50 mm2 and 20×20 mm2 were used in this research. The Al-doped SiOx
layers were formed on the Si surfaces through the processes described in Chapter 4; the treatments
were performed for 1 to 30 min. In addition, as a reference, after an RCA cleaning followed by
1% diluted hydrofluoric acid (HF) treatment for 1 min and rinse with DIW, these pieces were
immersed in the 70wt.% nitric acid (HNO3) aqueous solutions at room temperature (RT) for 10
min. Right after the HNO3 treatment and DIW rinse for 2 min, 80-nm-thick SiNy layers were142

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deposited on both sides or one side of samples by Cat-CVD at a substrate temperature of 100 °C,
SiH4 and NH3 flow rates of 4 and 250 sccm, respectively, under a pressure of 3.2 Pa. A tungsten(W) wire heated at 1800 °C was used as a catalyzer for decomposing those precursors into eachradical species. After the deposition for 700 s, the samples were placed in the quartz tube furnace
heated at 350 °C and annealed for 30 min in N2 gas under a flow rate of 1.0 L min-1 as the postdeposition annealing (PDA). In addition, the Cat-CVD SiNy layer was deposited on the fabricatedpyramidal textured Si substrate to demonstrate an excellent optical property at the same deposition
condition except for the processing time because it takes a longer time to deposit a same thicknessof ~80 nm on the pyramidal textured Si surface due to slopes with the base angle.

6.10.2 Sample evaluation
To evaluate the surface passivation, τeff and iVoc were analyzed as a function of Δn using a Sinton
lifetime tester (Sinton Instruments, WCT-120) for 50×50 mm2-sized pieces. Moreover, each Qf
was calculated from the C–V curves using the following equations [15] with obtained flat-band
voltage (VFB). For the calculation of Qf, the C–V curves have large hysteresis, which might berelated to the trapped charge in the Al-doped SiOx layer. Therefore, we determined VFB from theaverage of VFB values obtained by the forward and reverse scans.

𝑄f =

𝑉FB =

𝐶i
(𝛥𝛷ms − 𝑉FB )𝑞𝐴(𝑉FB,forward +𝑉FB,reverse )2

𝐶i =𝐴 𝜀r 𝜀0𝑑(6.12)(6.13)(6.14)

Where Ci is the insulator's capacitance, ΔΦms is the work function difference between metal andSi, A is the electrode area, and VFB is the flat band voltage. The electrode area, i.e., the contacted
mercury (Hg) probe area of 0.0040–0.0043 cm2, was estimated from the measured Ci, the insulator
thickness (d) of 80 nm, εr, and ε0 for each measurement.

143

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6.10.3 Surface passivation of SiNy/Al-doped SiOx stacks
Figure 6.27 (a) shows the τeff of samples treated with Al(NO3)3 concerning Δn at each treatment time.
The results indicate that τeff increases as treatment time increases, particularly in the low Δn region ofaround 1014 cm−3. This suggests that the Al-doped SiOx layer contains fixed charges. According toseveral papers [16–18], electron capture at the Al-induced acceptor states creates tetrahedrally four
coordinated Al oxide (AlO4−) with negative fixed charges. This AlO4− induces the inversion layer by
accumulating large holes at the Al-doped SiOx/Si interface, which leads to good surface passivation.
However, in the SiNy/Al-doped SiOx stacks configuration, significant degradation of τeff was observedin the low Δn region with increasing treatment time, as shown in Figure 6.27 (b). In this case, theopposite tendency can be attributed to the charge compensation of the positive fixed charges present
in the SiNy layers on top of the samples by the negative fixed charges in the Al-doped SiOx layers.

(b)τeff (µs)τeff (µs)(a)

Fig. 6.27 (a) τeff as a function of Δn for the Al(NO3)3-treated samples at the various treatment
time conditions. (b) τeff as a function of Δn for the samples with a single SiNy layer and SiNy/Aldoped SiOx stacks formed at various treatment times.

144

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ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

6.10.4 Fixed charge density of SiNy/Al-doped SiOx stacks
Figures 6.28 (a) and (b) depict the C–V curves for samples with single SiNy or SiNy/Al-doped
SiOx stacks formed at various Al(NO3)3 treatment times. The curves were acquired by sweeping
forward and reverse directions. The Al-doped SiOx layer significantly impacts the positive charge
reduction in the SiNy layer. This is evident from the C–V curves of the stacks, which are shifted
considerably to the positive voltage. Figure 6.29 shows the calculated Qf values as a function ofAl(NO3)3 treatment time. In the calculation, we used Eqs. (6.8) and (6.9), and the contacted Hgelectrode areas were estimated to be 0.0040–0.0043 cm2 from Eq. (6.10). The single SiNy layer
has a high positive Qf value of about 2×1012 cm−2. This value gradually decreases with increasing
the Al(NO3)3 treatment time in the SiNy/Al-doped SiOx stacks. Furthermore, negative Qf values
are confirmed for samples treated for more than 5 minutes. This implies that the Al-doped SiOx
layer is highly negatively charged and overcompensates the positive fixed charges in the SiNy
layer. The negative fixed charges can explain the surface passivation improvement anddegradation for the single Al-doped SiOx layer and the SiNy/Al-doped SiOx stacks shown in Figure
6.27 (a) and (b). The improvement of τeff for the single Al-doped SiOx layer with increasing
Al(NO3)3 treatment time is due to the increase of negative fixed charges in the layer. On the other
hand, the degradation of τeff for the SiNy/Al-doped SiOx stacks with increasing Al(NO3)3 treatment
time might be due to the charge compensation. Such a charge-compensated-related degradation
has also been reported. In the range of Δn as low as < 5×1013 cm−3, the passivation with a single
layer of Al2O3 showed a substantial decrease of τeff, which could be compensated by inserting an
HfO2 layer [19]. Moreover, in the case of Al2O3/SiO2 stacks, the negative Qf in Al2O3 wascompensated to ‘zero’ by ten cycles of ALD-SiO2 with the same order of magnitude of positive
Qf [19]. Therefore, the increased negative fixed charges overcompensate the positive fixedcharges localized in the SiNy layer, turning to the negative Qf from the positive Qf. Thus, theseexperimental results demonstrate the existence of high negative fixed charges in the Al-doped
SiOx layers. However, the estimated negative values of 1010–1011 cm−2 are insufficient to obtainthe surface passivation effect. SRV is rather increased around the low Qf values because holes still
exist as minority carriers near the n-Si surface, and carrier recombination actively occurs there.
Hence, the degradation mechanism can explain the extremely low τeff of the SiNy/Al-doped SiOx
stacks with low Qf values.

145

CHAPTER 6(a)

ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE(b)Fig. 6.28 (a) Forward- and (b) reverse-direction C–V curves for the samples with a single SiNy
layer and SiNy/Al-doped SiOx stacks prepared by Al(NO3)3 treatment for various treatment times.

Fig. 6.29

Qf for the samples with a single SiNy layer and SiNy/Al-doped SiOx stacks preparedby Al(NO3)3 treatment as a function of treatment time.
146

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ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

In addition, the effect of the non-doped SiOx interlayer underneath the SiNy was also
investigated. The τeff of SiNy/Al-doped SiOx and SiNy/non-doped SiOx as a function of Δn are shownin Figure 6.30. Aside from the significant τeff degradation for SiNy/Al-doped SiOx, a dramatic
improvement of τeff can be seen for SiNy/non-doped SiOx. Excellent τeff of 8.9 ms was confirmed,
corresponding iVoc of 730 mV and Seff, max of 1.6 cm s−1. Given that the oxides have a similar
thickness and SiO2 (Si4+) content, as shown in Figure 6.31 of fitted Si 2p core-level XPS spectra
and Table 6.4, charge compensation is likely responsible for the significant degradation of the τeff
in SiNy/Al-doped SiOx stacks and damage to the ultrathin layer is not the reason for the degradation.
To address this issue, a promising solution would be to deposit an ARC with a negative Qf, suchτeff (ms)as aluminum nitride (AlN) [20].

Fig. 6.30

τeff of SiNy/n-Si, SiNy/non-doped SiOx/n-Si, and SiNy/Al-doped SiOx/n-Si as a function of

Δn after 350 °C PDA.

147

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ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

(b)

(a)

Fig. 6.31

Si 2p core-level XPS spectra for (a) a SiNy/Al-doped SiOx stack and (b) a SiNy/non-doped SiOx stack.

Table 6.4

SiOx thickness and composition of the sub-oxides (Si1+, Si2+, and Si3+) and SiO2

(Si4+) components.
Thickness (nm)Si1+ (%)

Si2+ (%)

Si3+ (%)

Si4+ (%)

Al-doped SiOx

0.84

16.5

8.4

4.0

71.1

Non-doped SiOx

1.08

15.8

3.4

6.4

74.4

148

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ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

6.10.5 Optical properties of Cat-CVD SiNy on pyramidal textured Si
Figures 6.32 (a), (b), and 6.33 show reflectance, transmittance, and calculated absorbance spectraof Cat-CVD SiNy on pyramidal textured Si substrate, varying deposition time from 950 to 1100
s. The estimated JG values are listed in Table 6.5. The deposition time was optimized at 975 s to
obtain the highest JG value of 43.13 mA cm−2. This value demonstrated a significant increase in
the anti-reflective effect on Cat-CVD SiNy on pyramidal textured Si substrate, which realizes adrastic increase of carrier generation compared to other configurations of the front surface andexpected higher surface passivation quality by the Qf enhancement due to the increment of thegenerated carrier trapping. However, due to the charge compensation, SiNy on Al-doped SiOx
degrades the surface passivation quality. To enable excellent surface passivation, non-charged or
negative-charged ARC should be deposited.
(a)Fig. 6.32(b)(a) Reflectance and (b) transmittance spectra of Cat-CVD SiNy on planar andpyramidal textured Si substrate.

Fig. 6.33 Absorbance spectra of Cat-CVD SiNy on planar and pyramidal textured Si substrate.

149

CHAPTER 6
Table 6.5

ANTI-REFLECTIVE STRUCTURES AND LAYERS ON AL-COIL ARCHITECTURE

Photon current density (JG) for Cat-CVD SiNy on pyramidal textured Si substrate,

varying deposition time from 950 to 1100 s.
SiNy deposition time (s)JG (mA cm−2)

700

36.44

950

43.05

975

43.13

1000

43.06

1050

43.00

1100

42.94

SiNy on planar Si substrate

SiNy on pyramidal textured Si substrate

6.11 Summary
In this chapter, I investigated the effectiveness of antireflective structures and layers on Al-COIL
architecture. The findings revealed that the Al-COIL on a pyramidal textured Si substrate showeda higher surface passivation quality and photon current density, leading to improved solar cell
performance. The elevated surface passivation quality could be attributed to the improved lighttrapping effect. 1.54 times higher JG value implies that Qf might be enlarged by the electron
trapping in Al-induced acceptor states due to the increased carrier generation, causing an
improvement of surface passivation quality. Cat-CVD SiNy/ultrathin non-doped SiOx showed an
excellent surface passivation quality with a τeff of 8.9 ms corresponding to an iVoc of 730 mV and
optical properties with photon current density of >43 mA cm−2. Hence, Cat-CVD is a suitable
ARC deposition method for an ultrathin layer, as demonstrated experimentally. However, thedeposition of SiNy degraded the surface passivation of Al-COIL due to charge compensation.
Therefore, different ARCs without positive charges will be considered a promising solution. For
implementing device structure, the deposited ARC layer must be locally opened to make thecontact like a PERC structure. Hence, to remain Al-doped SiOx at the locally opened area, it isneeded to apply some complicated process such as laser ablation or masking and etching process
followed by further Al(NO3)3 treatment for the opened area. In order to pursue the low-cost
production, It should be realized by an alternative simple process in the future work.

150

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References[1]

K. R. McIntosh and S. C. Baker-Finch, "OPAL 2: Rapid optical simulation of silicon solar
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[2]

M. A. Green, "Lambertian light trapping in textured solar cells and light-emitting diodes:
analytical solutions," Prog. Photovolt. 10, 235 (2002).

[3]

K. R. McIntosh, and L. E. Black, “On effective surface recombination parameters,” J. Appl.
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C. T. Nguyen, K. Koyama, K. Higashimine, S. Terashima, C. Okamoto, S. Sugiyama, K.
Ohdaira, and H. Matsumura, “Novel chemical cleaning of textured crystalline silicon
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SiNx/amorphous silicon stacked layers,” Jpn. J. Appl. Phys. 56, 056502 (2017).

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G. von Gastrow, R. Alcubilla, P. Ortega, M. Yli-Koski, S. Conesa-Boj, A. F i Morral, H.
Savin, “Analysis of the Atomic Layer Deposited Al2O3 field-effect passivation in blacksilicon,” Sol. Ener. Mater. Sol. Cells 142, 29 (2015).

[6]

H. Savin, P. Repo, G. von Gastrow, P. Ortega, E. Calle, M. Garín, and R. Alcubilla, “Blacksilicon solar cells with interdigitated back-contacts achieve 22.1% efficiency,” Nat.
Nanotechnol. 10, 624 (2015).

[7]

S. Wang, X. Wu, F.-J. Ma, Y. Zhang, G. Scardera, D. Payne, M. Abbott, and B. Hoex,“Understanding field-effect passivation of black silicon: Modeling charge carrier
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Photovoltaic Specialists Conference, Online, pp. 590–594 (2020).

[8]

S. Wang, X. Wu, F.-J. Ma, D. Payne, M. Abbott, and B. Hoex, “Field-effect passivation ofundiffused black silicon surfaces,” IEEE J. Photovolt. 11, 897 (2021).

[9]J. J. Kopanski, M. Y. Afridi, S. Jeliazkov, W. Jiang, and T. R. Walker, “Scanning kelvinforce microscopy for characterizing nanostructures in atmosphere,” AIP Conf. Proc. 931,
530 (2007).

[10]

P. A. Basore, “3D numerical solutions inside the pyramid of textured silicon solar cells,”
AIP Conf. Proc. 2487, 040001 (2022).

[11]

J. Hölzl and F. K. Schulte, “Work Functions of Metals” in Solid Surface Physics, SpringerVerlag, Berlin (1979).

[12]

A. Goetzberger, J. Knobloch, and B. Voß, Crystalline Silicon Solar Cells (John Wiley &Sons, 1998).

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B. Mojrová, H. Chu, C. Peter, P. Preis, J. Lossen, V. D. Mihailetchi, and R. Kopecek, “Acomparison study of boron emitter passivation by silicon oxide and a PECVD silicon
nitride stack,” Energy Procedia 124, 288 (2017).

[14]

H. T. C. Tu, K. Koyama, C. T. Nguyen, K. Ohdaira, and H. Matsumura, “High-quality
surface passivation of crystalline silicon with chemical resistance and optical transparency
by using catalytic chemical vapor deposition SiNx layers and an ultrathin SiOx film,” Jpn.
J. Appl. Phys. 57 08RB17 (2018).

[15]

S. M. Sze, and Kwok K. Ng, Physics of Semiconductor Devices, 3rd edition(John Wiley & Sons, 2007).

[16]

D. Hiller, J. Göttlicher, R. Steininger, T. Huthwelker, J. Julin, F. Munnik, M. Wahl, W.
Bock, B. Schoenaers, A. Stesmans, and D. König, “Structural properties of Al−O
monolayers in SiO2 on silicon and the maximization of their negative fixed charge density”,
ACS Appl. Mater. Interfaces 10, 30495 (2018).

[17]

D. Hiller, P. M. Jordan, K Ding, M. Pomaska, T. Mikolajick, and D. König, “Deactivationof silicon surface states by Al-induced acceptor states from Al–O monolayers in SiO2,” J.
Appl. Phys. 125, 015301 (2019).

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B. Hoex, M. Bosman, N. Nandakumar, and W.M.M. Kessels, “C-Si surface passivation byaluminum oxide studied with electron energy loss spectroscopy,” In: Proceedings of IEEE
39th Photovoltaic Specialists Conference, Tampa, Florida, pp. 3333–3336 (2013).

[19]

D. K. Simon, Paul M. Jordan, I. Dirnstorfer, F. Benner, C. Richter, and T. Mikolajick,“Symmetrical Al2O3-based passivation layers for p- and n-type silicon,” Sol Ener. Mater.
Sol. Cells 131, 72 (2014).

[20]

G. Krugel, A. Sharma, W. Wolke, J. Rentsch, and R. Preu, “Study of hydrogenated AlN as
an anti-reflective coating and for the effective surface passivation of silicon,” Phys. Status
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152

CHAPTER 7

CONCLUSIONCHAPTER 7
CONCLUSION
In this study, novel hole-selective passivating contact was developed using a simple wet chemical
process. The study demonstrated that Al(NO3)3 treatment can form an ultrathin Al-doped SiOx
layer on Si substrates, which has an excellent surface passivation quality with hole-selectivity. On
a planar n-Si surface, the Seff, max reached as low as 16 cm s−1, corresponding to iVoc of 670 mV
and J0 of 65 fA cm−2. The hole selectivity (S10) value could be estimated to be 13.3 using the J0
and the analyzed ρc of 20 mΩ cm2, which is competitive to other selective contacts. XPS revealed
that the ultrathin Al-doped SiOx layer consists of a bilayer of AlOy/SiOx on Si substrates with an
Al surface atomic concentration of ~1 at.% and a SiOx thickness of ~1 nm. C–V measurement
revealed that intensely upward band bending at the ultrathin Al-doped SiOx is formed on the n-Si
surface.
The simplified SRH formalism and the analyzed Vbi estimated a Qf value of −3.21×1012 cm−2and a Nit value of 1.43×1012 cm−2. The estimated values indicate that the surface passivation
mechanism originated from field-effect and chemical passivation. In particular, the primary originof the surface passivation is attributed to the field effect passivation induced by a highly
negatively charged oxide, which provides strong upward band bending of the Si surface with anexcellent hole-selective contact property. The photoconductance measurement implies that thesurface passivation quality is enhanced by light illumination. The excited electron might betrapped in the Al-induced acceptor states existing in the Al-doped SiOx layer, forming negatively
charged AlO4− structures and enhancing field effect passivation induced by the fixed charges. The
light-induced surface passivation effectively occurred at the surface of pyramidal textured Si
154

CHAPTER 7

CONCLUSION

substrates. Even though the larger surface area leads to higher Nit on the pyramidal textured Si
surface, the surface passivation quality was higher than that of mirror-polished Si surfaces. The
Seff, max reached as low as 13.3 cm s−1, corresponding to iVoc of 682 mV and J0 of 59 fA cm−2. Byenhancing the excited electron trapping due to the light absorbance increase in the pyramidal
textured Si and the slight increase of the space charge compression in the structure, the lower Seffis realized. The ARC layer on the ultrathin Al-doped SiOx is also considered an initiator of surface
passivation, but the fixed charge in the ARC layer highly affects the surface passivation quality.
The uncharged or negatively charged ARC layer should be deposited on the ultrathin Al-doped
SiOx.
The light illumination also affects the Rsheet. Significant Rsheet reduction of two order magnitude
was confirmed for Si(111) surface as low as 317 Ω sq−1 under the 1 sun illumination. The enhanced
negative charges due to the enlargement of the electron trapping induce more surface hole density,
leading to the lower Rsheet. Hence, the light illumination activates the hole-selective passivating
contact property on the ultrathin Al-doped SiOx.
In addition, the solar cell performance demonstrated the device operation of Al-COIL solar
cells. The demonstration revealed that the induced inversion layer acts as a p+ layer on the n-type
Si substrates. However, the low performance in the η of 3.35%, especially in low Jsc and Voc values,
originated from the high surface recombination with the reduction of the Qf value as simulated.
Since the fine grid electrodes are applied on the emitter surface, and FF is improved, the low FF
originated from a still high sheet resistance. Furthermore, the pyramidal textured Si surface
improved all parameters of solar cell performance. The light-trapping enhances light absorbance,
resulting in a higher solar cell performance. However, the still low Voc might be due to surface
recombination caused by the reduction of Qf under the front side electrode. The light-illumination
effect cannot appear due to the shadow of the electrode and the low work function material of Ag
compared to the Si valence band edge.
To overcome this problem, applying a buffer layer or an Al-diffused layer will be a promisingsolution under the finger Ag electrode with a fine structure. In addition, to enhance the spacecharge region compression effect, nano-sized texture surface such as black Si should be tested asa surface structure underneath an ultrathin Al-doped SiOx layer, which possibly decreases the s
Rsheet of emitter surface, leading to the reduction of the front Ag area. Also, the novel solar cellthat fully exploits both n and p doping via electrostatic charges in an inversion layer solar cell will
be tested in the near future.

155

APPENDIX
Appendix 1: Simulation parameters for PC1DThe simulation parameters for the PC1D are shown in Figure A-1.1. The parameters were set toobtain the depth profile of carrier densities (Fig. 4.24) and the depth profile of energy bands (Fig.
4.25) for the negatively charged n-type Si surface.

Fig. A-1.1 Simulation parameters for the PC1D to obtain the depth profile of carrier densities
and the depth profile of energy bands for the negatively charged n-type Si surface.
155

The simulation parameters for the PC1D are shown in Figure A-1.2. The parameters were set
to obtain the photo J–V characteristics for the devices with surface charge-induced emittersassuming various Qf values in the range of −8.8×1010 – −1.0×1013 cm−2, enormous front-SRV of
1×1010 cm s−1 (Fig. 5.15) and without front-surface recombination (Fig. 5.16).

Fig. A-1.2

Simulation parameter for the PC1D to obtain the photo J–V characteristics for thedevices with surface charge-induced emitters.

156

The simulation parameters for the PC1D are shown in Figure A-1.3. The parameters were set
to obtain the depth profile of carrier densities of face-to-face structure with variable distances of
10 nm to 1 µm between both sides with highly negative Qf of 1×1013 cm−2 (Fig. 6.5). These results
provided information about the enhancement of carrier densities by a space charge region
compression.

Fig. A-1.3

Simulation parameters to obtain depth profile of carrier densities of face-to-facestructure.

157

Appendix 2: Simulation parameters for PC3S
The simulation parameters for the PC3S are shown in Figure A-2.1. The parameters were set toobtain a 2D mapping of surface hole density on (a) pyramidal textured and (b) planar Si surface
with a negative Qf of 1.0×1011 cm−3 (Fig. 6.21), 3.2×1012 cm−2, and 3.0×1013cm−3 (Fig. 6.22). The“SURFACE CHARGE” was variable, and “TEXTURE” was selected to “Pyramids” or “Planar.”
These results provided information about enhancing carrier densities by a space charge region
compression in tiny structures on pyramidal textured Si surfaces.

Fig. A-2.1 Simulation parameters to obtain a 2D mapping of surface hole density on pyramidal
textured Si with negative surface charges of 3.2×1012 cm−2.

158

Appendix 3: Simulation parameters for OPAL 2
The simulation parameter for the OPAL 2 to obtain generation rate (Fig. 6.2) and reflectance (Fig.
6.3). The “Surface morphology” was varied, Planar, upright random pyramids with a normal
angle of 54.74° or upright regular texture with a steeper angle of 80° were applied. The PECVD
SiNx (n=1.99) [1] was used from the data library as an ARC, and the thickness was optimized foreach surface morphology. One of the input parameters is shown in Figure A-3.1.

Fig. A-3.1

Simulation parameters to obtain generation rate and reflectance for the surface

morphology of upright random pyramids with a normal angle of 54.74° applied an optimized
PECVD SiNx (n=1.99).

159

Appendix 4: Photon flux spectrum of solar radiation
AM 1.5 G spectrum was used to calculate photon current density (JG) in each Si substrate witheach surface structure, such as planar and pyramidal texture. The used photon flux as a functionof wavelength [2] is shown in Figure A-4.1. The AM 1.5 G spectrum is a standard solar spectrum
designed for terrestrial flat plate modules and has an integrated power of 0.1 W cm−2.

Fig. A-4.1 Photon flux as a function of wavelength.

160

Appendix 5: Absorption coefficient of c-Si and penetration length in c-Si
The absorption coefficient (α) of c-Si [3] and penetration length in c-Si are shown in Figure A5.1 (a) and (b), respectively. The penetration length is defined as the decay length to be 1/eintensity. Hence, from Eq. (3.32), the penetration length can be described as 1/α. For the lightwith a wavelength of 350 nm, the penetration length would be ~0.0096 µm. On the other hand,for the light with a wavelength of 900 nm, the penetration length would be ~32.67 µm. Hence,
Short-wavelength light can absorb the surface of the c-Si substrate, but long-wavelength light canabsorb in bulk. Therefore, homogeneous generation requires a long wavelength of lightillumination. For the effective lifetime measurement, the light should be manipulated by cutting
short-wavelength or monochromatic laser light with a long wavelength. However, shortwavelength light should be used to achieve high generation at the surface region. In this study,
laser light with a wavelength of 349 nm was also used to activate the surface passivation.

(a)

(b)

Fig. A-5.1(a) Absorption coefficient (α) of c-Si and (b) penetration length in c-Si.

161

Appendix 6: Inelastic mean free path (IMFP)In the material, the inelastic collision of electrons occurs, and their energy will be lost. The
inelastic mean free path (IMFP) is the distance that decays 1/e signal intensity by inelastic
scattering. IMFP can be calculated using the TPP-2M equation [4] described below.

𝜆=𝐸
𝐸𝑝2 {𝛽 ln(𝛾𝐸) − 𝐶/𝐸 + 𝐷/𝐸 2(A-6.2)

𝐸𝑝 = 28.8√𝑁𝜈 𝜌/𝑀−0.5

𝛽 = −1.0 + 9.44(𝐸𝑝2 + 𝐸𝑔2 )(A-6.1)+ 0.69 𝜌0.1

(A-6.3)𝛾 = 0.191 𝜌−0.50(A-6.4)

𝐶 = 19.7 − 9.1 𝑈(A-6.5)

𝐷 = 534 − 208 𝑈(A-6.6)

𝐸𝑝2
𝑁𝜈 𝜌
=
𝑀
829.4

(A-6.7)

𝑈=

Where β, γ, C, and D are the parameters, E is the electron energy, Eg is the band gap energy, Nν
is the number of valence electrons, ρ is the density, and M is the atomic weight. IMFP can becalculated to be 3.07 nm for the Si using Eqs. (A-6.1) – (A-6.7) and below parameters. Hence,the escape depth of photoelectrons can be estimated to be 8.35 nm.

Table A-6.1 Parameters of the Si for IMFP calculation
Parameters

Units

ValuesEeV1387.5

Eg

eV

1.12

Nν

-

4

ρ

g cm−3

2.33

M

-

28.08

162

Appendix 7: τeff measurement using low excitation laser light photon density
Figure A-7 shows τeff measured using µ-PCD (PV-2000A, Semilab Co. Ltd.) with a low areal laser
photon density from 1.0×1011 to 1.5×1013 cm−2 and a wavelength of 904 nm. The τeff was gradually
decreased with increasing the areal photon density in the range of relatively low injection levels
between 1.0×1011 and 1.0×1012 cm−2. Since such significant improvement does not appear without
Si band bending, the high τeff at the lower injection level is attributable to field-effect passivation
due to the band bending of the Si surface. The generation of the minority carrier concentrationand the carrier recombination is significantly reduced by using an excitation laser of lower areal
photon density. Hence, this τeff dependence on carrier excitation density implies that the surface
passivation mechanism is ascribed to the field-effect passivation. However, the direction of theband bending cannot be determined by considering only this result. To elucidate the polarity of
the fixed charge density in the ultrathin Al-doped SiOx layer, C–V measurements were conducted,
as investigated in 6.10.4. However, τeff increased with the injection level from 1.0×1012 to 1.5×1013
cm−2. As shown in Figure 4.34(a) and described in 4.7.4, this increase might be attributable to theenhancement of the negative Qf induced by electron trapping in the Al-induced acceptor states
due to the instantly increased Δn by the illumination of the pulsed laser light, increasing the
upward band bending and enhancement of field-effect passivation. The µ-PCD measurements
were performed at the Fukushima Renewable Energy Institute, AIST.

Fig. A-7.

τeff measured using µ-PCD with a low areal laser photon density from 1.0×1011 to 1.5×1013

cm−2 and a wavelength of 904 nm.
163

Appendix 8: Surface passivation by different aluminum precursor
To reveal the effectiveness of Al(NO3)3 aqueous solution to form ultrathin Al-doped SiOx layers on cSi substrates, the different aluminum precursor was tested to determine whether they could realize theexcellent surface passivation of the c-Si surface. In this study, aluminum sulfide (Al2(SO4)3) was used
for the Al doping and oxidant. The aluminum sulfide hydroxide (Al2(SO4)316H2O) was dissolved in
DIW to make 0.5 M Al2(SO4)3 aqueous solution, and the boiled Al2(SO4)3 aqueous solution was used.
HF-treated n-Si substrates were immersed in the boiled Al2(SO4)3 aqueous solution for 3 to 30 min.
After the DIW rinse followed by N2 blow dry, τeff measurements were performed for each sample.
Figure A-8 shows the τeff of Al2(SO4)3-treated samples for each process time, as well as the τeff ofAl(NO3)3-treated samples. As shown in Figure A-8, apart from Al(NO3)3-treated samples, the τeff of
Al2(SO4)3-treated samples remained low, around 13 to 18 µs in the entire range of the processing time.
As an observation of the hydrophobicity on the Al-treated c-Si surface, hydrophobic surfaces werestill observed for all of the samples. This implies that sufficient oxidation did not occur for theAl2(SO4)3-treated samples, leading to low surface passivation.

Fig. A-8.

τeff of Al2(SO4)3- and Al(NO3)3-treated samples for each process time.

164

Appendix 9: Stability and recovery of surface passivation quality
The stability of the ultrathin Al-doped SiOx layer on n-Si(100) was investigated. Figure A-9-1 shows
the time-related τeff transition through 30 days of the samples after Al(NO3)3 treatment with a boilingtemperature and concentration of 0.25 M for 5–30 min. 5 min treated sample was stabilized at a low
τeff. However, the samples with relatively high initial τeff showed gradual degradation with passing theelapsed time after the surface passivation. A peculiar increase feature was confirmed on a 10 mintreated sample. However, this sample also showed a gradual degradation after 3 days later. The lackof initial surface passivation quality was compensated with further oxidation in the ambient air,resulting in higher surface passivation. Another crucial finding was the τeff recovery by the lightillumination. Figure A-9-2 (a) shows τeff mapping after the time-related degradation, and (b) shows τeff
mapping after the 1 sun illumination for 5 min on the degraded sample. 1 sun illumination not fully
but recovered at a certain level of surface passivation quality with a τeff of 393 µs, as shown in Table
A-9. This result implies that the surface passivation induced by the ultrathin Al-doped SiOx layer canretain the surface passivation quality even in the ambient air but under light illumination (carrier
excitation). The possible reason for the τeff recovery is Qf recovery. The slight reduction from the initialτeff might be attributable to the further interfacial SiOx growth, leading to the mitigation of electron
traps in the Al-induced acceptor states owing to the less tunneling probability through the thickened
SiOx layer during the storing in the air ambient.

Fig. A-9-1.

Time-related τeff transition through 30 days of the samples after Al(NO3)3 treatment

with a boiling temperature and concentration of 0.25 M for 5–30 min.

165

(a)

Fig. A-9-2.

(b)(a) τeff mapping after the time-related degradation, (b) τeff mapping after the 1 sunillumination for 5 min on the degraded sample.

Table A-9 Mean and maximum τeff of after the degradation and after the 1 sun illumination
Substrate

Mean τeff

Maximum τeff

After degradation

311

416

After 1 sun illumination

393

491

166

Appendix 10: Surface passivation quality on p-type c-Si
Surface passivation quality on p-type c-Si (p-Si) was also investigated as performed on n-Si in
this study. The same process has proceeded as described in 4.3.1. The ultrathin Al-doped SiOx
layers were formed on the RCA-cleaned and HF-treated p-Si(100) substrates by immersing in the
boiled Al(NO3)3 aqueous solution with a molar concentration of 0.25 M for several processing
time conditions in the range of 5–30 min. Figure A-10 shows the τeff of the Al(NO3)3-treatedsamples on p-Si(100) as a function of the Al(NO3)3 processing time. Up to 15 min, τeff improved
with increasing the processing time. However, after 15 min, τeff gradually decreased withincreasing the processing time. The optimized time condition of 15 min realized 406 µs as a meanτeff. The tendency of the increment and decrease with respect to the processing time can be
explained as described in 4.4.3. Considering the induction of negative Qf originated from AlO4
structures formed by trapping the excited electrons in Al-induced acceptor states, more adjacent
to stoichiometric SiO2 stabilizes to create this AlO4 structure with increasing the processing time.
However, further oxidation makes further oxide growth at the interface, which suppresses theelectron tunneling through the oxide and trapping in the Al-induced acceptor states. Hence, furtheroxide growth inhibits the negative Qf induction, causing the τeff degradation.

Fig. A-10.

τeff of the Al(NO3)3-treated samples on p-Si(100) as a function of the Al(NO3)3

processing time.

167

Appendix 11: Frequency dependence of C–V measurements
The C–V characteristics of the ultrathin Al-doped SiOx/n-Si interface were investigated. To mitigate aleakage current, ~80 nm-thick SiNy layers were deposited as capping layers on the Al-doped SiOx layer.
The sample fabrication process and evaluation method are shown in 6.10.1 and 6.10.2, respectively.
During the C–V measurements, the frequency was varied in the range of 100 Hz to 1 MHz. Figure A11 (a) shows C–V curves measured with the frequency of 100 Hz to 1 MHz for 1 min Al(NO3)3-treatedsamples and Figure A-11 (b) for 12.5 min Al(NO3)3-treated samples. As shown in Figure A-11 (a),inversion capacitances were not observed for a 1 min Al(NO3)3-treated sample. On the other hand,inversion capacitances were observed for a 12.5 min Al(NO3)3-treated sample, as shown in Figure A11 (b). Furthermore, the inversion capacitances were significantly increased with decreasing
measurement frequency. Considering negative Qf induction in a 12.5 min-treated sample but still
positive Qf induction in a 1 min-treated sample, these results suggest that negative Qf induces aninversion layer at the surface of the n-Si substrate. The inversion capacitances can be observed byelectron trapping at the inversion condition by applying a negative bias to the top electrodes. The lower
frequency measurements increase inversion capacitance owing to the increment of the interface trap
response. Hence, a sufficient Al(NO3)3-treated sample creates an inversion layer induced by highly
negatively charged Al-doped SiOx with electron-trapped Al-induced acceptor states.

(a)

(b)

Fig. A-11. C–V curves (a) for 1 min Al(NO3)3-treated sample and (b) for 12.5 min Al(NO3)3-treatedsample measured with the frequency of 100 Hz to 1 MHz.

168

Appendix 12: τeff degradation by residual water
The Al(NO3)3 treatment increases τeff to ~550 µs at one of the best conditions (0.25 M, Boilingtemperature, 12.5 min) acquired by µ-PCD. However, sometimes, such a high τeff could not be obtained
just after the surface passivation. Figure A-12-1 shows τeff mapping just after the Al-doped SiOx
formation. As shown in Figure A-12-1, the τeff was one order lower than successfully well-passivatedthe Si surface at the same condition of Al(NO3)3 treatment. The data was acquired just after the DIW
rinse, followed by the N2 blow. The N2 dry process might not completely remove the water moleculesfrom the sample surface. To investigate the reason for the low τeff, the confirmed well-passivatedsample was immersed in the DIW in a transparent plastic bag and measured τeff using µ-PCD. Figure
A-12-2 shows the τeff (a) after N2 dried and (b) during immersing in DIW. These results suggest thatwater significantly affects the τeff degradation: ~300 µs was decreased to ~10 µs. The experimental
fact implies that one of the reasons for the τeff degradation is residual water on the surface passivation
layer. Considering that the oxidation-reduction potential of H2O (H2O + e− → 1/2H2 + OH−) is −0.83
eV, the trapped electrons in the Al-induced acceptor states occur a reduction reaction of the residual
water, resulting in reducing the upward band bending.

Fig. A-12-1

τeff mapping just after the Al-doped SiOx formation. The N2 dry process might notcompletely remove the water molecules from the sample surface.

169

(b)

(a)

Fig. A-12-2

τeff of Al(NO3)3-treated n-Si(100) (a) after N2 dried and (b) during immersing in DIW.

In addition, the τeff during immersing in ethanol was investigated. Since the sample immersed inethanol did not show significant τeff degradation as a mean τeff of 431 µs, the ethanol could replace theresidual H2O, mitigating the reduction reaction of H2O and the τeff degradation. Hence, the rinse
process with ethanol improves the issue of τeff degradation due to the residual water.

170

Appendix 13: Fabrication of inverted pyramidal texture
The inverted pyramidal-textured c-Si substrates were formed using a wet chemical process by referringto the reported literature [5]. The etching solution was prepared (the mixture of 0.04 M copper (II)nitrate trihydrate Cu(NO3)3, 4.6 M HF, 0.55 M H2O2, and 2 mM polyethylene glycol (PEG) with anaverage molecular weight of 200) and then immersed in the etching solution for 15 min at 50 °C withstirring at 1000 rpm. The experimental setup for fabricating the inverted pyramidal textured c-Si
substrates is shown in Figure A-13-1. The residual Cu precipitates were removed by immersing in 70
wt.% HNO3 aqueous solution for 20 min with ultrasonication. In this study, the effectiveness of PEG
was investigated. Figure A-13-2 shows the reflectance spectra of the formed inverted pyramidal
texture by immersing in the etching solution w/ or w/o PEG. In the wide range of wavelengths, thereflectance was slightly decreased. At a wavelength of 630 nm, the w/ PEG sample showed a lower
reflectance of 3.0% compared with the w/o PEG sample of 4.2%. These reflectance values were
exceptionally low compared with the conventional upright pyramidal texture with a reflectance of
7.8% at a wavelength of 630 nm. As shown in the SEM images of Figure A-13-3, the size of theinverted pyramidal texture is almost the same, but the w/ PEG sample surface seems smoother thanthat of the w/o PEG sample. Further surface etching on the inverted pyramidal texture surface using
NH4OH aqueous solution was necessary to obtain an excellent surface passivation quality, as describedin Appendix 14.

Fig. A-13-1. Experimental setup for the fabrication of the inverted pyramidal textured c-Si substrates.

171

Fig. A-13-2.

Reflectance spectra of the formed inverted pyramidal texture by immersing in the

etching solution w/ or w/o PEG.

Fig. A-13-3.

SEM image of the formed inverted pyramidal texture by immersing in the etchingsolution w/ or w/o PEG.
However, the additional surface etching increased reflectance, as shown in Figure A-13-3. To obtainlower reflectance compared with the conventional upright pyramidal texture, the surface etchingprocess time should be less than 2 min. Further etching is considered to destroy the shape of theinverted pyramidal texture by the isotopic alkaline etching, deteriorating the optical property withhigher reflectance.
172

Fig. A-13-3.

Reflectance spectra of the formed inverted pyramidal texture by immersing in the

etching solution w/ or w/o additional etching using NH4OH aqueous solution.

173

Appendix 14: Surface passivation on inverted pyramidal texture
The etching solution was prepared (the mixture of 0.04 M copper (II) nitrate trihydrate Cu(NO3)3, 4.6
M HF, 0.55 M H2O2) and then immersed in the etching solution for 15 min at 50 °C with stirring at
1000 rpm. After the residual Cu precipitates removal by immersing in 70 wt.% HNO3 aqueous solution
for 20 min with ultrasonication followed by oxide removal using a mixture of 40% CH3OH and 5%
HF for 5 min, the slight surface etching was performed using 28% NH4OH for 0.5–2 min at RT. Afteroxide removal by immersion in a 5% HF aqueous solution, SC-2 cleaning was conducted at 80 °C for
10 min. After 5% HF treatment, ultrathin Al-doped SiOx was formed by immersing in 0.25 M Al(NO3)3
aqueous solution at boiling temperature for 12.5 min. As shown in Figure 14-1 and Table A-14, the τeff
was gradually increased with increasing the additional surface etching time using NH4OH aqueous
solution. Hence, additional surface etching on the inverted pyramidal texture was essential to obtain
excellent surface passivation quality. However, as displayed in the surface images, the blackness isslightly reduced with the increase in additional etching time. This is due to the reflectance change as
described in Appendix 13 and shown in Figure 13-3.

w/o NH4OH

w/ NH4OH

w/ NH4OH

w/ NH4OH

etching

etching (0.5 min)etching (1 min)etching (2 min)Surfaceimageτeffmapping

Fig.A-14-1. Surface images and τeff mappings after the surface passivation by Al-doped SiOx of theformed inverted pyramidal textured c-Si w/ or w/o additional etching using NH4OH aqueous solution.

174

Table A-14

τeff of w/ and w/o additional etching of inverted pyramidal texture using NH4OH

w/o NH4OH

w/ NH4OH

w/ NH4OH

w/ NH4OH

etching

etching (0.5 min)etching (1 min)etching (2 min)Mean τeff (µs)

6

70

83

189

Maximum τeff (µs)

14

257

287

392

Substrate

PEG was added to the etching solution to form an inverted pyramidal texture, as described in Appendix
13. After performing the same cleaning processes, ultrathin Al-doped SiOx was formed on the inverted
pyramidal texture by immersing in 0.25 M Al(NO3)3 aqueous solution at boiling temperature for 20
min. The mean τeff was 300 µs, and the maximum τeff was 421 µs. Hence, this result demonstrated an
excellent surface passivation quality can be realized on inverted pyramidal textured c-Si surfaces.
However, the additional etching process was essential to obtain excellent surface passivation quality,
but it increased reflectance. Therefore, this trade-off between surface passivation quality andreflectance must be overcome to realize higher properties at the same time.

Fig. A-14-2. τeff mapping after the surface passivation by Al-doped SiOx of the inverted pyramidal
textured c-Si formed in PEG added etching solution and w/ additional etching for 2 min using NH4OH
aqueous solution.

175

Appendix 15: PEDOT:PSS on ultrathin Al-doped SiOx
Poly(3,4-ethylenedioxythiophene) polystyrene sulfonate (PEDOT:PSS) is known as a hole transportlayer with high conductivity and high work function. By using PEDOT:PSS dispersion solution, a thinfilm can be formed on a substrate surface. As described in 4.7.3, the Rsheet of Al-doped SiOx/n-Si is stillhigh to obtain an excellent solar cell performance with a sufficient FF. One of the promising solutions
for the issue is utilizing a high-hole conductive layer on the ultrathin Al-doped SiOx layer. Thus, toovercome this issue, PEDOT:PSS on the ultrathin Al-doped SiOx layer was investigated to determine
whether it can realize a low Rsheet while keeping an excellent surface passivation quality. In thisexperiment, 280 µm-thick FZ-grown n-Si(100) substrate with a resistivity of 2.5 Ωcm was used. Aftercleaning in acetone followed by ethanol solvent with ultrasonication for 5 min, respectively, and 1%
HF treatment for 1 min at RT, ultrathin Al-doped SiOx was formed by immersing in 0.25 M Al(NO3)3
aqueous solution at boiling temperature for 12.5 min. Subsequently, PEDOT:PSS (Heraeus
Deutschland GmbH & Co. KG, Clevios PH1000) was spin-coated on the ultrathin Al-doped SiOx/nSi(100) using spin coater (MIKASA Co., Ltd., 1H-DX) following the process, as shown in Figure A15-1. Figure A-15-2 shows τeff after the spin coating of PEDOT:PSS at various spin speeds (2000–
5000 rpm). Since the high τeff of 420 µs before the spin coating significantly degraded to less than 90µs, PEDOT:PSS layer degraded the surface passivation of ultrathin Al-doped SiOx. The reason is not
fully revealed but the acid of PEDOT:PSS might destroy the ultrathin Al-doped SiOx layer,
deteriorating the surface passivation. As shown in Figure A-15-3, Rsheet was gradually decreased with
reducing spin speed due to the thickening of the PEDOT:PSS layer. However, a sufficient low Rsheet
could not obtained. Hence, this experiment concluded different approaches should be investigated toincrease the lateral conductivity of the ultrathin Al-doped SiOx layer/n-Si.

Fig. A-15-1. Spin coating process flow to form PEDOT:PSS thin film on the Al-doped SiOx.

176

Fig. A-15-2.

τeff after the spin coating of PEDOT:PSS at various spin speeds (2000–5000 rpm).

Fig. A-15-3. Rsheet and thickness of PEDOT:PSS layer formed at various spin speeds (2000–5000
rpm).
177

Appendix 16: Al/Mg/ultrathin Al-doped SiOx stacks
As shown in Figure 6.24, the Al-COIL solar cell was tested as a rear structure of Al/Mg/ultrathin Aldoped SiOx stacks. As described in this thesis, since the Al-doped SiOx layer is revealed to be a holeselective contact, the rear Al-doped SiOx layer must be removed or transformed the layer into electronselective contact. The removal of Al-doped SiOx makes direct metal contact with the Si surface,resulting in high recombination at the rear surface. Hence, a better solution is the transformation of
the carrier selectivity of the Al-doped SiOx layer. In order to change the carrier selectivity, we shouldmake contact with a low-work function metal. Magnesium (Mg) is one of the low-work function
metals (3.66 eV) with respect to an electron affinity in Si of 4.05 eV. The concept is shown in Figure
A-16; when the low-work function metal is in contact with the surface of Al-doped SiOx/n-Si, the Si
surface band is bent downwards by creating an accumulation layer. Hence, the Al-induced acceptor
states transform into the donor states which can provide electron-selective contacts.

Fig. A-16-1. Band diagram before and after Mg contact with ultrathin Al-doped SiOx/n-Si

The tunneling current density of the Al/Mg/Al-doped SiOx/n-Si structure was evaluated. The ultrathin
Al-doped SiOx layers were formed on the RCA-cleaned and HF-treated n-Si(100) substrates by
immersing in the boiled Al(NO3)3 aqueous solution with a molar concentration of 0.25 M for 12.5 min.
A Mg wire with a diameter of 0.5 mm (Nilaco Corporation, 99.95%) was used to deposit Mg. After178

the Mg electrode formation, Al was subsequently thermally deposited on the Mg/Al-doped SiOx/n-Si
through the mask with a diameter of 0.0028 cm2 to form circular Al/Mg stack electrodes without
breaking the vacuum condition with a pressure of < 9×10−4 Pa. After the deposition, the oxide of therear side was removed using 5% HF and by scratching. Subsequently, full-area Al film was depositedon the rear side of the sample to form an Ohmic contact. To compare the Al/Mg stacks and only Al asa top electrode, the Al/Al-doped SiOx/n-Si/Al structure was also fabricated and tested. Figure A-16-2
shows dark J–V characteristics of the samples with the MOS structures with Al/Mg/Al-doped SiOx /nSi/Al (w/ Mg) and Al/Al-doped SiOx/n-Si/Al (w/o Mg). The J–V characteristics of the w/ Mg sample
indicate an increase in the tunneling current density with more Ohmic properties compared with the
w/o Mg sample. This implies that Mg contact transforms the Al-doped SiOx/n-Si into electronselective contact owing to the downward band bending of the Si surface due to the low-work functionof Mg.

Fig. A-16-2. Dark J–V characteristics and the device structure of the samples with Al/Mg/Al-doped
SiOx /n-Si/Al (w/ Mg) and Al/Al-doped SiOx/n-Si/Al (w/o Mg).

179

Appendix 17: Surface passivation on rough Si surface
The effectiveness of the rough Si surface on surface passivation by ultrathin Al-doped SiOx wasinvestigated. An acid etching process was performed to make a rough surface of Si substrates, asreported in the literature [6]. A 280 µm-thick n-Si(100) wafer with a resistivity of 2.5 Ωcm was
immersed in the acid mixture of 96.67% HCl, 0.83% HF, and 2.5% H2O2 at RT for 60 min to form arough Si surface. After the etching process, the treated substrates were cleaned using acetone andethanol with an ultrasonication. Subsequently, 1% HF treatment was performed for 1 min. After DIW
rinse, the hydrophobic Si substrates were immersed in the boiled Al(NO3)3 aqueous solution with amolar concentration of 0.25 M for 12.5 min. After the N2 blow, τeff values were measured by µ-PCD
using an excitation laser with a wavelength of 904 nm and an areal photon density of 5×1013 cm−2. By
the acid etching process, the Si substrates were 10 µm thinned. Hence, the measured τeff values wereconverted into Seff, max to compare the surface passivation quality on the mirror-polished Si surface. As
shown in Figure A-17 of the optical microscope image of the c-Si surface after the acid etching process,~10 µm-size bubble-like textures were observed.

Fig. A-17. Optical microscope image of the c-Si surface after the acid etching process.

180

Table A-17 shows the average mean τeff and Seff, max of the Al(NO3)3-treated sample with and without
the bubble-like textures. For each sample, four samples were tested and evaluated as their averagevalues. The bubble-like textures increase the surface defect states corresponding to the surface area of
Si. Nevertheless, the Seff, max was slightly decreased compared with the mirror-polished Si surface.
Hence, the surface roughness would be an advantage of the surface passivation by the ultrathin Aldoped SiOx. One possible reason is that effective negative Qf is increased owing to the increment of
the surface area, enhancing the field-effect passivation.

Table A-17

average mean τeff and Seff, max of the Al(NO3)3-treated sample with and without the

bubble-like textures

Substrate

Average mean τeff

Average mean Seff, max

w/ bubble-like textures

550

25.5

w/o bubble-like textures

476

30.2

181

Appendix 18: PMMA on Al-doped SiOx for anti-reflection coating
The ultrathin Al-doped SiOx cannot functionalize as an anti-reflection coating. Hence, an antireflection coating is essential to boost photocurrent density. In general, the PECVD-SiNy layer or
indium tin oxide (ITO) is widely used. However, the plasma or the sputtering causes the degradationof the ultrathin Al-doped SiOx layer, leading to poor interface properties. In this thesis, the Cat-CVD
SiNy layer was tested for the anti-reflection coating on Al-doped SiOx, but the surface passivation was
significantly degraded, as described in 6.10.3. The degradation is considered to be due to Qf
compensation by the high positive Qf in the Cat-CVD SiNy layer. Therefore, to avoid such damage tothe ultrathin layer and Qf compensation, a wet deposition process using a material with a refractive
index of ~2 increases optical properties. According to the optical simulation results using OPAL 2, asshown in Figure A-18-1, poly(methyl 2-methyl propanoate) (PMMA) is one of the promising materials.
PMMA expects the reflected photon current density loss to be as low as 1.16 mA cm−2 when the
thickness was optimized as 109 nm. In this experiment, PMMA (Sigma-Aldrich GmbH, 182230) was
dissolved in acetone to prepare 2 wt.% PMMA solution. The PMMA dispersive solution was spincoated on the ultrathin Al-doped SiOx/n-Si(100) with a spin speed of 500 rpm for 5 s followed by 2500
rpm for 30 s, as illustrated in Figure A-18-2. The τeff was measured by µ-PCD before and after spincoating. The τeff was 443 µs before pain-coating but degraded to 53 µs after spin-coating of PMMA.
The τeff degradation mechanism has not been revealed yet, but acrylic resin might be related to thedegradation because attaching scotch tape consisting of acrylic resin also caused degradation as low
as 8 µs, but attaching cello tape without acrylic resin did not cause such degradation. As shown in
Figure A-18-3, the reflectance of PMMA-coated Al-doped SiOx/Si was lower than that of the sample
without PMMA layer. However, the anti-reflection effect is insufficient due to the still unoptimized
thickness. Although the reflectance can be optimized by adjusting the concentration of the PMMA
solution and spin speed during the spin coating process, the τeff degradation issue cannot be solved aslong as it is not replaced with other materials.

182

Fig. A-18-1. Optical simulation results using OPAL 2 for the structure with PMMA/SiO2/Si.

183

Fig. A-18-2. Spin coating process for the formation of PMMA anti-reflection coating.

Fig. A-18-3. Reflectance spectra of PMMA-coated Al-doped SiOx/Si and Al-doped SiOx/Si.

184

Appendix 19: WOz on Al-doped SiOx for anti-reflection coating
As described in Appendix 18, forming an anti-reflection coating by wet deposition process is crucial
to realize higher Al-COIL solar cell performance due to the boosted photocurrent density. In addition,
high-work function materials can be hole-transport layers owing to the band alignment by receiving
electrons from the Si surface and enhancing the upward band bending of the Si surface with strong
inversion layers. Furthermore, the enhanced carrier generation by the anti-reflection coating is
considered to increase the excited electron trapping in the Al-induced acceptor states, resulting inenhancing negative Qf in the Al-doped SiOx followed by field-effect passivation and hole-lateral
conductivity. The concept of the enhancement of hole-selective passivating contact properties
originating from boosting the light trapping effect is illustrated in Figure A-19-1.

Fig. A-19-1. Concept of the enhancement of hole-selective passivating contact properties originatingfrom boosting the light trapping effect.

For the material with anti-reflection and high-work function properties, WOz was selected because thematerial has an adequate refractive index of 2.45 and a work function of ~6.7 eV. The opticalsimulation was conducted using OPAL2 to analyze the loss due to the reflectance. WOz expects thereflected photon current density loss as low as 1.41 mA cm−2 when the thickness was optimized as 56
nm.

185

Fig. A-19-2. Optical simulation results using OPAL 2 for the structure with WO3/SiO2/Si.
In this experiment, cubic WOz (c-WOz) nanoparticles dispersed in a xylene solution were used. The
WOz dispersed solution was spin-coated on ultrathin Al-doped SiOx/n-Si(100). Subsequently, thereflectance spectra and τeff were measured using UV-vis NIR and µ-PCD, respectively. Figure A-19-3
shows the reflectance spectra of spin-coated c-WOz/Al-doped SiOx/n-Si. Consequently, c-WOz
effectively reduced reflectance with respect to the bare-Si surface, and the bluish surface can be seen186

in the inset of the figure. The lowest reflectance was still a slightly longer wavelength compared withthe optimal wavelength. The thickness and refractive index should be adjusted by investigating thespin coating conditions to obtain higher optical properties. Figure A-19-4 shows τeff after the depositionof the WOz layer measured by µ-PCD at each measurement time to investigate the time-related changeof τeff. Although the τeff degraded just after the deposition, τeff gradually recovered to the initial value
while storing in the air ambient at RT but slightly degraded in 24 h and stabilized at ~390 µs. However,higher τeff could not be realized by depositing WOz on Al-doped SiOx/n-Si. This concluded that the
expectation of the high τeff owing to the higher band bending due to the high-work function of c-WOz
formed by the spin-coating process could not be realized. The spin-coated c-WO might not have ahigh-work function, resulting in reducing the upward band bending. In addition, spectroscopic
ellipsometry measurement and analysis using the Cauchy model revealed the c-WOz layer thickness
was 100 nm, and the refractive index at 630 nm was 1.66. The refractive index was not far from thetheoretical value of 2.45 for WO3 crystallite.

Fig. A-19-3. Reflectance spectra of spin-coated c-WOz/Al-doped SiOx/n-Si. As references, thereflectance spectra of mirror-polished Si, Cat-CVD SiNx, and simulated amorphous WO3 are also
displayed.

187

Before depo.

After depo.

After depo.

After depo.

After depo.

3 h later

24 h later

48 h later

435

393

387

τeff
mappingMeanτeff (µs)

453

356

Fig. A-19-4. τeff after the deposition of c-WOz measured by µ-PCD at each measurement time.
In addition, WOz was deposited using an electron beam evaporation system on ultrathin Al-doped SiOx
for both sides. The thickness of WOz was determined to be ~36 nm by spectroscopic ellipsometry
measurement and analysis using the Cathey model. As a result, the deposition of WOz significantly
degraded the τeff of n-Si(100) passivated by Al-doped SiOx from ~400 µs to 15 µs. Since the C–V
measurement revealed the Vbi reduction from 0.87 to < 0.8 eV, the deterioration is mainly due to the
reduction of the upward band bending. The work function of deposited WOz was not so higher thanthe expected value of ~6.7 eV. Therefore, WOz/Al-doped SiOx stacks could not be a solution for the
excellent hole-selective passivation contact on n-Si substrates. However, the surface passivation
quality may depend on the thickness of WOz and may need to be optimized.

188

Appendix 20: Fabrication of nano-sized pyramidal textured c-Si surface
Space charge region compression in the narrow n-Si with the faced negatively charged oxides
increases hole density. Figure A-20-1 shows the PC3S simulated surface hole density increase ratio of
the pyramidal texture surface with respect to the planar surface as a function of the texture depth. The
Qf of −3×1012 cm−2, the texture angle of 54.7°, and the doping concentration of base Si of 2×1015 cm−3
without impurity diffusion concentration were set and simulated in the region of the base thickness of
15 µm. As shown in Figure A-20-1, the surface hole density is gradually enhanced by decreasing the
texture depth. The texture size of 0.5 µm makes the surface hole density increase ratio with respect tothe planar surface > 1.4 times higher. Since the surface recombination velocity and sheet resistanceare inversely proportional to the surface hole density, forming nano-sized pyramidal texture is one of
the solutions to improve hole selective passivating contact properties. Hence, in this study, nanosized pyramidal textures were fabricated using a conventional alkaline wet chemical etching processwith additives as an etching mask to enhance the hole-selective passivating contact properties.

Fig. A-20-1.

Surface hole density increase ratio with respect to planar surface as a function of

texture depth.

189

In this experiment, 280 µm-thick n-Si(100) with a resistivity of 2.5 Ω cm was used as substrates.
After cleaning the cleaved substrates in acetone followed by ethanol with ultrasonication for 5 min,the substrates were immersed in the heated KOH solution with additives (TP101, Hayashi Pure
Chemical) to form pyramidal textures on the Si substrates. The details of the texturing conditions areshown in Table A-20.

Table A-20. Conditions for the formation of nano-sized pyramidal textures

Conditions

Temperature

Time

Additives

Ultrasonication

(°C)

(min)

(wt.%)

Ref.

80

15

0.3

-

Low temp.

55

45

0.3

-

Low temp. w/ US

55

45

0.3

AppliedLow temp., ×2 additives w/ US

55

45

0.6

AppliedLow temp.,×4 additives w/ US

55

45

1.2

AppliedFigure A-20-2 shows the SEM images of the fabricated pyramidal textures for each condition. Twomain factors were process temperature and additives concentration to fabricate the nano-sizedpyramidal textures. As compared to the ref. (a) and (b) decreasing the process temperature from 80 to
55 °C significantly decreased the pyramidal texture size. The ultrasonication effect on the texture size
was also investigated, but almost no effect was confirmed, as compared to (b) and (c). However, afurther increase of additives concentration from 0.3 to 1.2 wt.% decreased the pyramidal texture size
from ~1 µm to ~500 nm. Hence, the significant reduction of the pyramidal texture size from ~2 µm to~500 nm was caused by lower process temperature and increasing additives concentration. However,the decreasing pyramidal textures slightly increased the reflectance in the wide range of wavelengths
except for the short wavelength region of the smallest pyramidal textures, as shown in Figure A-20-3.
This increase of the reflectance might be related to decreasing texture depth. Assuming this model, thehigher aspect ratio of nano-sized pyramidal textures is considered to be a solution to realize both nanosize structure and lower reflectance.

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(a)

(b)

(c)

(d)

(e)

Fig. A-20-2.

SEM images of the surface of the fabricated pyramidal textures for each condition. (a)reference, (b) 55 °C condition, (c) 55 °C w/ ultrasonication, (d) 55 °C, 0.6 wt.% additives w/
ultrasonication, and (e) 55 °C, 1.2 wt.% additives w/ ultrasonication.

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Fig. A-20-3.

Reflectance spectra of the fabricated pyramidal textures for each condition.

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Appendix 21: Potassium nitrate (KNO3) treatment on c-Si surface
As described in this thesis, the Al(NO3)3 treatment boosts the surface passivation quality on c-Si
surfaces. Other metal nitrates may oxidize the Si surface, and the included metal acts as a field-effectpassivation as a formation of metal oxide with a low or high work function compared to an electron
affinity of c-Si or metal ion with a positive charge. Hence, in this experiment, potassium nitrate(KNO3) was investigated to determine whether it could improve the surface passivation quality on thec-Si surface. Cleaved 280 µm-thick FZ n-Si(100) substrates with a resistivity of 2.5 Ω cm were cleaned
using RCA cleaning and treated in 1% HF solution. The hydrophobic c-Si substrates were immersed
in 0.5 mol L−1 KNO3 aqueous solution at RT for 8 h. The τeff mapping of Figure A-21-1 (a) shows thata slightly increased spot was observed up to ~8 µs. Hence, since the KNO3 treatment passivates the Si
surface, and the surface passivation quality surpassed the passivated surface by HNO3 (τeff ~5 µs), thesurface passivation might be attributable to the field-effect passivation owing to the included metal
potassium as a formation of ion. Moreover, to improve the surface passivation quality by KNO3
treatment, the KNO3 concentration condition of 2.0 mol L−1 for 2 h increased to τeff of 28 µs and using
2.0 mol L−1 KNO3 aqueous solution prepared by SC-1 solution for 10 min further increased to 111 µs,
as shown in Figure A-21-2. The positive ion of K+ might be attracted and attached by the negatively
charged Si or SiOx surface in the alkaline solution, resulting in the induction of field-effect passivation
owing to the higher positive K+ surface. The proposed surface passivation model is shown in Figure
A-21-3.

(a)Fig. A-21-1.

(b)τeff mapping of (a) 0.5 mol L−1 KNO3 treated sample and (b) 70 wt.% HNO3 treatedsample.

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(a)

(b)

Fig. A-21-2.

τeff mapping of (a) 2.0 mol L−1 KNO3 treated sample and (b) 2.0 mol L−1 KNO3 in SC-

1 solution treated sample.

Fig. A-21-3.

Surface passivation model by formed ultrathin K+ doped SiOx layer formed by KNO3

treatment.

194

References[1]

S. Duttagupta, F. Ma, B. Hoex, T. Mueller, and A. G. Aberle, “Optimised antireflection
coatings using silicon nitride on textured silicon surfaces based on measurements andmultidimensional modelling,” Energy Procedia 15, 78 (2012).

[2]

C. A. Gueymard, "SMARTS2, A simple model of the atmospheric radiative transfer ofsunshine: Algorithms and performance assessment," Florida Solar Energy Center Report(1995).

[3]

M. A. Green, “Self-consistent optical parameters of intrinsic silicon at 300 K including
temperature coefficients,” Sol. Energy Mater. Sol. Cells 92, 1305 (2008).

[4]

S. Tanuma, C. J. Powell, and D. R. Penn, “Calculations of electron inelastic mean freepaths. V. Data for 14 organic compounds over the 50–2000 eV range,” Surf. Interface Anal.
21, 165 (1994).

[5]

Y. Wang, L. Yang, Y. Liu, Z. Mei, W. Chen, J. Li, H. Liang, A. Kuznetsov, and D. Xiaolong,“Maskless inverted pyramid texturization of silicon,” Sci. Rep. 5, 10843 (2015).

[6]

A. Stapf, F. Honeit, C. Gondek, and E. Kroke, “Texturing of monocrystalline silicon wafers
by HF-HCl-H2O2 mixtures: Generation of random inverted pyramids and simulation oflight trapping in PERC solar cells,” Sol. Energy Mater. Sol. Cells 159, 112 (2017).

195

LIST OF PUBLICATION
Journal articles[1]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Surface orientation
dependence on hole-selective passivating contacts of ultrathin Al-doped SiOx layers formedon n-type crystalline Si surfaces using an aluminum nitrate oxidation treatment,” Appl.
Surf. Sci. (Submitted)[2]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Effectiveness of the
Pyramidal Textured Crystalline Silicon Surfaces on Hole-Selective Passivating Contacts
Induced by Wet Chemically Grown Ultrathin Al-doped SiOx layers,” Solar RRL(Submitted)[3]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Significant role of an
ultrathin SiOx interlayer in Cat-CVD SiNy/SiOx stacks for the surface passivation of n-type
crystalline silicon,” Prog. Photovolt. (Submitted)[4]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Ultrathin Al-doped SiOx
passivating hole-selective contacts formed by a simple wet process,” Jpn. J. Appl. Phys.
62, SK1040 (2023).

[5]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Hole-Selective Ultrathin AlDoped SiOx Passivation Layer Formed by Immersing in Aluminum Nitrate Aqueous
Solution,” Phys. Status Solidi (RRL) 16, 2200052 (2022).

International Conference[1]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Effectiveness of Al-induced
charged oxide inversion layer on pyramidal textured Si surfaces,” 34th International
Photovoltaic Science and Engineering Conference (PVSEC-34), Shenzhen, China (6–10th
Nov. 2023) (Poster presentation)[2]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Excellent Surface
Passivation on a Pyramidal Textured Si Surface by Wet Chemically Grown Al-induced
Charged Oxide Inversion Layer,” 40th European Photovoltaic Solar Energy Conference andExhibition (EU PVSEC), Lisbon, Portugal (18–22nd Sep. 2023) (Oral presentation)
196

[3]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “The significant role of an
ultrathin SiOx interlayer underneath a Cat-CVD SiNy passivation layer for the highefficiency n-type crystalline Si solar cells,” 11th International Conference on Hot-Wire
Chemical Vapor Deposition (11th HWCVD), Jülich, Germany (17–19th Apr. 2023) (Oralpresentation)[4]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Ultrathin Al-doped SiOx
passivating hole-selective contacts formed by a simple wet process,” 33rd International
Photovoltaic Science and Engineering Conference (PVSEC-33), Nagoya, Japan (13–17th
Nov. 2022) (Oral presentation)[5]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Cat-CVD SiNx /ultrathin
SiOx passivation stacks for high-efficiency crystalline Si solar cells,” 31st International
Photovoltaic Science and Engineering Conference (PVSEC-31), Online (13–15th Dec.
2021) (Poster presentation)

Domestic Conference[1]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “The effectiveness of the p+
inversion layer on an n-type crystalline Si surface with a pyramidal texture formed byaluminum nitrate oxidation treatment,” The 3rd Japan Photovoltaic Society Conference,
Kyoto, Japan (29–30th Jun. 2023) (Oral presentation)[2]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Surface passivation ofpyramidal textured Si surfaces by Cat-CVD SiNy/ultrathin SiOx stack layers,” 20th CatCVD research meeting, Kitakyushu, Japan (22–23rd Jun. 2023) (Poster presentation)[3]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “The application of theinversion layers induced by the ultrathin Al-doped SiOx layers for the front-emitter n-type
crystalline Si solar cells,” The 70th JSAP Spring Meeting 2023, Tokyo, Japan (15–18th Mar.
2023) (Oral presentation)[4]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Surface passivation
mechanism on Si surfaces by Al-doped ultrathin SiOx formed by immersing in aluminum
nitrate aqueous solution,” The 83th JSAP Autumn Meeting 2022, Sendai, Japan (20–23rd
Sep. 2022) (Oral presentation)[5]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Effectiveness of Cat-CVD
on Al-doped ultrathin SiOx layers,” 19th Cat-CVD research meeting, Tosu, Japan (14–15th
Jul. 2022) (Poster presentation)[6]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Orientation dependence ofSi surface passivation for an ultrathin Al-doped SiOx layer,” The 2nd Japan Photovoltaic
Society Conference, Kanazawa, Japan (28–29th Jun. 2022) (Poster presentation)
197

[7]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Surface passivation on Si
surfaces by Al-doped ultrathin SiOx formed by wet process,” The 69th JSAP Spring Meeting
2022, Kanagawa, Japan (22–26th Mar. 2022) (Oral presentation)[8]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Surface passivation
mechanism in Cat-CVD SiNx/ultra-thin SiOx stacks,” JSAP Hokuriku-Shinetsu Branch
Meeting 2021, Shinshu, Japan (4th Dec. 2021) (Oral presentation)[9]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Surface passivation on
crystalline silicon surfaces by Cat-CVD SiNx/NAOS-SiOx stacks,” The 1st Japan
Photovoltaic Society Conference, Online (14–15th Oct. 2021) (Oral presentation)[10]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Surface passivation on
crystalline silicon surfaces by SiNx/SiOx stacks formed by catalytic chemical vapor
deposition and nitric acid oxidation,” The 82nd JSAP Autumn Meeting 2021, Online (10–13th Sep. 2021) (Oral presentation)[11]

Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “Effect of ultra-thin siliconoxide film formed at the Cat-CVD SiNx/Si(100) interface on surface passivation,” 18th CatCVD research meeting, Nomi, Japan (9–10th Jul. 2021) (Poster presentation)

Award
Innovative Photovoltaic Young Researcher’s Award
Hiroki Nakajima, Huynh Thi Cam Tu, and Keisuke Ohdaira, “The effectiveness of the p+ inversion
layer on an n-type crystalline Si surface with a pyramidal texture formed by aluminum nitrateoxidation treatment,” The 3rd Japan Photovoltaic Society Conference, Kyoto, Japan (29–30th Jun.
2023) (Oral presentation)

Grant
JST SPRING: Support for Pioneering Research Initiated by the Next Generation, Grant Number
JPMJSP2102

198