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University of Wollongong University of Wollongong
Research Online Research Online
University of Wollongong Thesis Collection 2017+ University of Wollongong Thesis Collections
2018
Surface and Interface Engineering in Semiconducting Electrocatalyst and Surface and Interface Engineering in Semiconducting Electrocatalyst and
Photo(electro)catalyst Photo(electro)catalyst
Haifeng Feng University of Wollongong
Follow this and additional works at: https://ro.uow.edu.au/theses1
University of Wollongong University of Wollongong
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Recommended Citation Recommended Citation Feng, Haifeng, Surface and Interface Engineering in Semiconducting Electrocatalyst and Photo(electro)catalyst, Doctor of Philosophy thesis, Institute for Superconducting and Electronic Materials, University of Wollongong, 2018. https://ro.uow.edu.au/theses1/430
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Surface and Interface Engineering in Semiconducting
Electrocatalyst and Photo(electro)catalyst
This thesis is presented as part of the requirements for the
Award of the Degree of
Doctor of Philosophy
from the
University of Wollongong
by
HAIFENG FENG
B. Sc., M. Sc.
Supervisors:
Dr. Yi Du, Dr. Xun Xu, Prof. Shi Xue Dou
Institute for Superconducting and Electronic Materials
Faculty of Engineering
July 2018
i
Table of contents
Table of contents........................................................................................................... i
Acknowledgements ..................................................................................................... iv
Certification ................................................................................................................. v
List of abbreviations................................................................................................... vi
Abstract ....................................................................................................................... ix
Chapter 1 Introduction ............................................................................................... 1
1.1 General background ............................................................................................ 1
1.2 Literature review and research motivation .......................................................... 2
1.2.1 Semiconductor electrocatalysts .................................................................... 2
1.2.2 TiO2 based catalysts ..................................................................................... 4
1.2.3 Semiconductor photocatalyst and photo(electro)catalysts ......................... 11
1.2.4 BiOBr based catalysts ................................................................................ 14
1.3 Outline of chapters ............................................................................................ 17
1.4 References ......................................................................................................... 18
Chapter 2 Experimental Techniques ....................................................................... 32
2.1 Scanning tunneling microscopy ........................................................................ 32
2.1.1 Quantum tunneling ..................................................................................... 32
2.1.2 Working principle of STM ......................................................................... 34
2.1.3 Ultra-high vacuum system and molecular beam epitaxy ........................... 36
2.2 Atomic force microscopy .................................................................................. 38
2.2.1 Working principles of AFM ....................................................................... 38
2.2.2 Working modes of AFM ............................................................................ 39
2.3 Other facilities ................................................................................................... 43
2.3.1 Scanning electron microscopy and focused ion beam ............................... 43
2.3.2 Transmission electron microscopy ............................................................. 44
2.3.3 Photoelectron spectroscopy ........................................................................ 47
2.4 References ......................................................................................................... 49
Chapter 3 Characterization of TiO2(110) surface by STM ................................... 51
3.1 Introduction ....................................................................................................... 51
3.2 Experimental section ......................................................................................... 52
3.3 Results and discussion ....................................................................................... 53
3.3.1 Rutile(110) surface with (1×1) reconstruction ........................................... 53
3.3.2 OV point defects on (1×1) surface ............................................................. 54
3.3.3 Rutile(110) surface with (1 × 2) reconstruction ......................................... 56
ii
3.4 Summary ........................................................................................................... 58
3.5 References ......................................................................................................... 58
Chapter 4 Activating titania for electrocatalysis by surface vacancy engineering
..................................................................................................................................... 62
4.1 Introduction ....................................................................................................... 62
4.2 Experimental sections ....................................................................................... 63
4.3 Characterization of OV by STEM ..................................................................... 65
4.4 Characterization the electronic structure by in-situ synchrotron XPS and UPS68
4.5 Electrochemical HER activity of OV-TiO2 ....................................................... 71
4.6 Determination of the origin of the electrochemical HER activity of OV-TiO2 73
4.6.1 Electrical conductivity................................................................................ 73
4.6.2 Electrochemical activate sites .................................................................... 75
4.6.3 DFT calculation of the Gibbs free energy .................................................. 80
4.7 Summary ........................................................................................................... 87
4.8 References ......................................................................................................... 87
Chapter 5 Tuning electronic structure of BiOBr 2D nanosheets by strain for
photocatalysis ............................................................................................................. 93
5.1 Introduction ........................................................................................................... 93
5.2 Experimental section ......................................................................................... 94
5.3 Morphology and structure characterization ....................................................... 96
5.4 Characterization of strain in BiOBr 2D nanosheets .......................................... 97
5.5 Characterization of photocatalytic activity of BiOBr 2D nanosheets ............. 100
5.6 DFT of the strain effect on the electronic structure of BiOBr ........................ 102
5.7 Summary ......................................................................................................... 105
5.8 References ....................................................................................................... 105
Chapter 6 Construction of 2D lateral pseudo-heterostructures by strain
engineering ............................................................................................................... 109
6.1 Introduction ..................................................................................................... 109
6.2 Experimental section ....................................................................................... 111
6.3 Structure characterization of BiOBr nanosheets ............................................. 113
6.4 Strain induced pseudo-heterostructure ............................................................ 116
6.5 Electronic properties characterization ............................................................. 122
6.6 DFT calculation of the strain effect on the band structure of BiOBr .............. 126
6.7 Performance of BiOBr nanosheets in photoelectronic devices ....................... 128
6.8 Summary ......................................................................................................... 132
6.9 References ....................................................................................................... 132
Chapter 7 Summary and outlooks ......................................................................... 137
iii
Appendix A: List of publications .............................................................................. 139
Appendix B: Conference contributions ..................................................................... 141
iv
Acknowledgements
This thesis was accomplished under the supervision of Dr. Yi Du, Dr. Xun Xu, and
Prof. Shi Xue Dou in Institute for Superconducting and Electronic Materials (ISEM)
in University of Wollongong.
Firstly, I want to acknowledge Dr. Du, my principle supervisor, for both his strong
supports in experiments and building my scientific thoughts. I am also grateful to my
co-supervisors, Dr. Xu and Prof. Dou, for their enthusiastic guidance and significant
contributions during the whole PhD course. Their meticulous scientific attitudes and
creative ways of thinking will be a wealth that benefit to my entire working career.
I would also like to thank all the colleagues, also my friends, in our research group.
I received continuous help from Dr. Jincheng Zhuang, Dr. Yundan Liu, Miss Yani Liu
and Dr. Zhi Li in STM operation, data analysis and daily discussions. Meanwhile, I am
grateful for other group members for their kindness in internal collaborations and
discussions, including Mr. Amar Al-Keisy, Miss Li Wang, Mr. Long Ren, Mr. Liang
Wang, Mrs. Nana Wang, Dr. Zhongchao Bai, Miss, Ningyan Cheng, Miss Nana Liu
and Miss Guyue Bo. Thanks Assoc. Prof. Michael Higgins and Dr. Tian Zheng from
IPRI for AFM training. Kind helps from Mr. Robert Morgan, Mr. Paul Hammersley,
Mr. John Wilton, Mr. Mat Davies for maintaining STM equipment are appreciated.
In addition, I need to thank my collaborators, Prof. Weichang Hao, Dr. Zhongfei Xu,
and Mrs. Dandan Cui in Beihang University and UOW-BUAA Joint Research Centre,
Prof. Jun Chen in IPRI, Prof. Zhenpeng Hu in Nankai University, Prof. Jijun Zhao and
Mr. Nan Gao in Dalian University of Technology.
At last, I want to take this opportunity of thanks my parents and my wife for their
endless love and support during this four years.
v
Certification
I, Haifeng Feng, declare that this thesis, submitted in fulfilment of the requirements for
the award of Doctor of Philosophy, in the Institute for Superconducting & Electronic
Materials, Faculty of Engineering, University of Wollongong, is wholly my own work
unless otherwise referenced or acknowledged. This document has not been submitted
for qualifications at any other academic institution.
vi
List of abbreviations
HER hydrogen evolution reaction
OER oxygen evolution reaction
LDOS local density of states
TMO transition metal oxide
XAS x-ray absorption spectroscopy
STM scanning tunneling microscopy
STS scanning tunneling spectroscopy
DFT density functional theory
AFM atomic force microscopy
KPFM kelvin probe force microscopy
SKPM scanning kelvin probe microscopy
CPD contact potential difference
UHV ultra-high vacuum
TMP turbo molecular pump
TSP titanium sublimation pump
MBE molecular beam epitaxy
LT low temperature
GV gate valve
SEM scanning electron microscopy
FIB focused ion beam
EDX energy-dispersive x-ray spectroscopy
TEM transmission electron microscopy
vii
SAED selected area electron diffraction
HRTEM high resolution transmission electron microscopy
STEM scanning transmission electron microscopy
ABF annular bright-field
HAADF high angle annular dark-field
EELS electron energy loss spectroscopy
GPA geometric phase analysis
XPS x-ray photoelectron spectroscopy
UPS ultraviolet photoelectron spectroscopy
UV ultraviolet
UV-vis ultraviolet-visible
CB conduction band
VB valence band
CBM conduction band maximum
VBM valence band maximum
OV oxygen vacancy
OHP hydroxyl group pairs
CNT carbon nanotube
MO methyl orange
Rh B rhodamine B
CTAB cetyl trimethylammonium bromide
FFT fast Fourier transform
XRD x-ray diffraction
PPMS physical property measurement system
VASP vienna Ab initio simulation package
viii
GGA generalized gradient approximation
PBE Perdew-Burke-Ernzerhof
PAW projector augmented wave
ZPE zero point energies
LSV linear sweep voltammetric
EASA electrochemically active surface area
RHE eversible hydrogen electrode
ix
Abstract
Surface and interface engineering is one of the most effective approaches in tuning
semiconductors for their chemistry (energy, environment and catalysis) and device
applications (electronic devices, and optical-electronic devices).
In this thesis, we show several approaches of modifying the catalytic and electronic
properties of several semiconductors through manipulating their surface and interface.
Techniques including scanning probe microscopies (SPM), electron microscopies, and
electron spectroscopies, and, X-ray spectroscopies, were taken used to characterize the
effect of surface and interface engineering, and the effect on their electronic properties.
This thesis includes:
1. Oxygen vacancies (OV) engineering on TiO2 rutile(110) single crystal. By carefully
controlling the annealing temperature and annealing time, OV defects can be precisely
introduced into the single crystal. The evolution of the surface structure was
investigated by in-situ low temperature STM. It was found that at 900 K, OV point
defects isolating with each other exist on the (1 × 1) surface. In addition, OVs on the
surface tended to be filled either water molecule or OHs. When the annealed
temperature increased to 1300 K, (1 × 2) surface reconstruction dominated on the
surface, in which cross-links defects caused by more oxygen deficiency were found.
2. Activating TiO2 for electrocatalysis by surface vacancy engineering. STEM
techniques clarified that the good crystalline nature of reduced TiO2 with and the spatial
distribution of OV in the surface and near-surface region. A mid-gap defect electronic
state was created by the OVs, which can effective enhance the electric conductivity of
the reduced TiO2. The reduced TiO2 exhibit tremendous enhancement in hydrogen
x
evolution reaction (HER) due to the increased electric conductivity and amounts of
active sites. Combing the in-situ observation of the water splitting reaction by STM
and DFT calculations, subsurface oxygen vacancies and low coordinated Ti ions (Ti3+)
demonstrated their roles in enhancing the electrical conductivity and promoting
electron transfer and hydrogen desorption, which activate reduced TiO2 single crystal
in the hydrogen evolution reaction in alkaline media. This study offers a rational route
for developing reduced transition metal oxide for low-cost and highly active hydrogen
evolution reaction catalysts, to realize over all water splitting in alkaline media.
3. The band structure of BiOBr 2D nanosheets was tuned by the strain for
photocatalysis. The inner strain in the BiOBr nanosheets has been tuned continuously
by controlled manipulating their shapes. The photocatalytic performance of BiOBr in
dye degradation can be manipulated by the strain effect. The low-strain BiOBr
nanosheets show improved photocatalytic activity. DFT suggest that strain can modify
the band structure and symmetry in BiOBr. The enhanced photocatalytic activity in
low-strain BiOBr nanosheets is due to improved charge separation attributable to a
highly dispersive band structure with an indirect band gap.
4. Constructing two dimensional (2D) lateral pseudo-heterointerface by strain
engineering in BiOBr nanosheets. Taking advantage of their strain-sensitive layer
structure, 2D lateral pseudo-heterogeneous interface are realized in the single-
component BiOBr nanosheets by finely tune the pH value in synthesis. Due to the
proper band alignment cross the interface, charge separation under visible light
irradiation was enhanced, which was reflected by the photo-current measurements and
the degradation experiments of pollutions. The strain engineering was demonstrated to
xi
be an effective way to tune the electronic structure of BiOBr and promote its efficiency
in photoenergy conversion applications. In addition, the construction of the lateral
pseudo-heterostructure through strain exhibit promising applications in building
unprecedented 2D systems with exciting properties.
1
Chapter 1 Introduction
1.1 General background
Semiconducting electronic materials can be defined as the semiconductors which are
applied for their electrical properties. According to band theory, a semiconductor
material with small but non-zero band gap, has an electric conductivity between that of
an insulator and that of most metals. The behaviours of electrons in semiconductors are
determined by their band structure, which has been successfully implemented in
describing many physical properties of solids. They have been utilized in a variety of
application fields mainly based on their different electronic properties, including
chemistry (energy conversion and storage, environmental remediation and catalysis)
and device applications (electronic devices, and optical-electronic devices).[1-2]
Surface and interface plays a critical role in determining the electronic properties of
semiconductor materials, especially for those in nanostructures with very large specific
surface area.[3-5] Abundant surface mismatches, low coordinated ions or atoms, and
surface defects can strongly alter their electronic structures, which could play a vital
role in the corresponding performances in their applications.[6-8] However, effectively
characterizing conditions of surface or interface, and modifying the electronic
properties through controllably manipulating surface conditions are challenging, due
to the lack of approaches in monitoring fine surface structure and localized variation of
electronic structure in micro or nano-scale. Effectively control the species,
concentration, and distributions of surface defects and their effects on the electronic
structure, hence, are of great significance in developing their applications.
2
1.2 Literature review and research motivation
1.2.1 Semiconductor electrocatalysts
In the research filed of energy conversion, including fuel cells, solar cells and
batteries, electrocatalysts are crucial in the reaction rate, efficiency, and selectivity.[9]
Pursuing efficient and low-cost electrocatalysts is one of the key topics in this filed. In
this thesis, we will largely limit the introduction to water electrolysis that split water
into hydrogen and oxygen. For both the hydrogen evolution reaction (HER) and oxygen
evolution reaction (OER), noble metal (Pt, Ir, Au) or noble metal based electrocatalysts
with high activity and stability are still dominating in these applications, which are not
desirable due to their high cost and low-abundance.
Varieties of earth-abundant semiconductor materials have been extensively explored
as alternatives of noble metal based electrocatalysts, in which the most efficient ones
are designed based on transition metal elements (Fe, Co, Ni, Cu, Mo, W, Mn, Ti, Zn)
and p-block elements (S, Se, C, O, N, P).[10-16] Among them, transitional metal oxides
(TMOs), which are being widely applied either as catalysts or supports in industrial
processes, have been recognized as appealing alternatives to noble-metal based
electrocatalysts for the OER, especially in alkaline media.[17-19]
To reach a high electrocatalytic activity close to noble metals, two requirements are
needed to be met in the first place. The first one is a high conductivity, which enable
the transport of electrons between electrodes and active sites. To overcome the poor
conductivity of common semiconductor materials, conductive materials, for examples
carbon based materials, such as, carbon based materials and Ni foams, are need to be
used as conducting substrate to fabricate a heterogeneous structure.[20-22]
3
Figure 1.1 (a) Trend in overpotential for the OER is shown as a function of the 3d
transition elements. (b) Trend in overpotential for the HER is shown as a function of
the 3d transition elements. In both reactions, the trends of Mn < Fe < Co < Ni was
determined.[23] Reproduced with permission from Springer Nature.
The second important aspect is the active site on the surface of an electrocatalyst, in
which the catalytic charge transfer happens. Identifying the active site of high active
electrocatalysts are important in further improving their activity, revealing the catalytic
mechanism and rationally developing new electrocatalysts. For example, for an early
work shown in Figure 1.1, combined with X-ray absorption spectroscopy (XAS),
4
scanning tunneling microscopy (STM) and electrochemical measurements,
Subbaraman et al. proposed that the strength of OH-M2+δ (0 ≤ δ ≤ 1.5) energetic (M,
Ni < Co < Fe < Mn) play the determining role in the reactivity of both OER and HER
in alkaline media (Mn < Fe < Co < Ni), which was helpful for guiding further designing
electrocatalysts by tuning active sites.[23]
Inspired by these works, intense attempts have been made in developing TMOs for
HER and overall water splitting in recent years. Nanoscale NiO/Ni-carbon nanotube
(CNT) material[24], Ni/CeO2−CNT[25], 3D crystalline/amorphous Co/Co3O4 core/shell
nanostructure on Ni foam[26], oxygen vacancy (OV)-CoO/carbon[27] and TiO2
nanodots/Co nanotubes on carbon fibers[28] have exhibited competitive HER activities
toward commercial Pt/C catalysts in alkaline media. Hollow Co3O4 microtube arrays
on nickel skeleton,[29] porous MoO2 nickel foam,[30] Ni−Co−A (A = P, Se, O)
nanosheets on nickel foam[31] and have been reported to be highly active bifunctional
electrocatalysts for overall water splitting.
1.2.2 TiO2 based catalysts
Figure 1.2 The crystal structure of the three phases of TiO2. (a) anatase (tetragonal, a =
0.3785 nm, c = 0.9513 nm), (b) rutile (tetragonal, a = 0.4593 nm, c = 0.2959 nm), (c)
brookite (orthorhombic, a = 0.5455 nm, b = 0.9181 nm, c = 0.5142 nm).
TiO2 is found exist in nature in three crystal structures.[38] The crystal structure of
the three phases is shown in Figure 1.2. Rutile, a tetragonal structure, is the most
5
common one in nature with the most stable structure. Due this reason, rutile is the most
studied oxide surface although in the two decades.[32] Beside rutile, anatase is also
tetragonal structure, while brookite is orthorhombic. In recent years, the study of
anatase in fields of both surface characterization and surface catalysis have also got
huge progresses.[33-37]
In the study of catalytic properties of TiO2, rutile(110) face is mostly selected,
because it is the most thermodynamically stable crystal face of TiO2. The information
and knowledge we got from this surface are also useful for the understanding of other
crystal faces and phases of TiO2, as well as other oxide surfaces. As shown in Figure
1.2 (a), rutile TiO2 presents a tetragonal structure, with one Ti atom surrounded by six
O atoms distributed in a distorted octahedral. In the bulk structure, Ti atoms are six-
fold (6f) and oxygen atoms are threefold coordinated. While, on the rutile(110) surface
as shown in Figure 1.3, there are two types Ti atoms, five-fold (5f) and 6f coordinated
Ti atoms, as well as two-fold (2f) coordinated oxygen atoms (bridging oxygen atoms)
and three-fold (3f) coordinated oxygen atoms. On the perfect stoichiometric rutile TiO2
(110) surface, 5f Ti atom row, 6f atom rows and 3f oxygen atom rows lie in a plane of
along the [001] direction, while bridging oxygen atom rows locate above 6f Ti rows in
the same direction. The surface of (110) cleaved rutile has been proved to be auto-
compensated and charge neutral, as well as breaks the minimum number of bonds and
the longest bonds in the crystal structure, which consequently is the lowest-energy
surface structure, by means like first-principles calculations and STM.[39,40] For the
freshly cleaved surface, the first few layers on the TiO2 rutile(110) structure will relax
to saturate the dangling bonds. As a result, a rumpling structure is created by two
different types of Ti (5f Ti and 6f Ti) atoms moving in opposite direction on the top
several atomic layers. The unit cell of TiO2 rutile(110)-(1×1), therefore, is defined as
6
two adjacent 5f Ti atoms in [001] direction is with a distance 2.96 Å, and two adjacent
5f Ti atoms in [1-10] direction with a distance of 6.5 Å.[41-44]
Figure 1.3 Diagram of the crystal structure of TiO2 rutile(110) surface, with 5f Ti, Obr
and surface OV marked.[44]
Figure 1.4 STM images of TiO2(110)-(1×1) surfaces with different surface species. (a)
a reduced surface with OVs, (b) partially hydroxylated surface with OH groups, and
(c) fully hydroxylated surface with OHPs. (d) Height profiles of the labeled sites in (b)
for OV, OH, and OHP, respectively. All the STM images have a size of 11 nm × 11
nm, and were obtained at V = 1.6 V, I = 10 pA. [45] Reprinted with permission from
American Chemical Society.
7
Plenty of STM images have been acquired from TiO2 rutile(110)-(1×1) surface
which are consistent with the theoretical structure in Figure 1.3. Most images are
obtained on empty-state mode, because normally on filled-state mode the tip will
scrape on the surface. Limited images on filled-state mode were achieved by some
groups.[46,47] The bright rows along [001] direction are composed of 5f Ti atoms, while
the dark rows are assigned to bridging oxygen atoms on empty-state imaging mode.
This contrast is dominated by the distribution of density of states despite bridging
oxygen atoms are higher than 5f Ti atoms in physical geometry structure.
Figure 1.5 STS measurements on TiO2 rutile(110)-(1×1) surfaces. (a) dI/dV of the
occupied state of different site at 78 K: a, at the center of a lobe in the occupied STM
image between the two 5f Ti sites indicated in the inset STM image. b, at a 6f Ti site.
c, at the OVbr site. Insets are the STM image and the I-V curves corresponding to the
three dI/dV curves.[47] Reproduced with permission from AIP Publishing. (b) STS
measured above OVbr and above regular 5f Ti surface atoms on TiO2 rutile(110) surface
at 78 K, with the measurement range from -3 V (occupied state) to 1.5 V (empty
state).[48] Reproduced with permission from American Physical Society.
Many types of surface defects have been observed on TiO2 rutile(110)-(1×1)
surface. Three types of intrinsic point defects in the reduced TiO2 rutile(110), including
OV defect marked as I, hydroxyl group (OH group) marked as II and hydroxyl group
8
pairs (OHP), are shown in Figure 1.4. Normally, in ultra-high vacuum (UHV) OVs are
created by the reduction of rutile by heating or ion sputtering, while OH groups are
created by the dissociation of water at the oxygen vacancies. Although these defects
are all bright by STM and locate at the middle of two adjacent 5f Ti atoms rows, the
heights at same scanning conditions by STM are different. For example, as shown in
Figure 1.4, the height of OHP (1.6 V, 10 pA) is around 1.2 Å and 1.0 Å for OH at the
same scanning conditions. While for OV the height is much less than OH and OHP. It
is also found that the relative apparent height of OH groups are sensitive to the tip-
sample distance in STM measurements, which were not observed on OV and 5f Ti
atoms.[45]
Figure 1.6 (a)-(d) Dynamic processes of the dissociation of water at an OV site on
TiO2(110) at about 187 K. Protrusions are labeled as follows: OV, OHbr groups and
water on 5f Ti sites are marked as open white circles, filled white circles and filled
black squares, respectively. (e) Schematic diagram of the dynamic processes of water
dissociation on the rutile(110) surface.[50] Reproduced with permission from American
Physical Society.
9
Besides getting the surface morphology with atomic resolution of TiO2, the
electronic band structure TiO2 has been studied by STM, which is strong relevant to its
performance in many applications.[45,46,48] As shown in Figure 1.5 (a), Minato et al. use
scanning tunneling spectroscopy (STS) to identify the occupied Ti 3d defect state of
TiO2. While in the work of Setvin et al. (Figure 1.5 (b)) no clear differences between
the LODS of 5f Ti atoms and OV were observed.
Figure 1.7 Two sets, (a)-(b) and (c)-(e) show two reaction processes between Oa and
water at 300 K, with the interpretation illustrated by the diagram below.[57] Reproduced
with permission from American Physical Society
10
Since the Fujishima and Honda found the photocatalytic water splitting by TiO2, the
study of the interaction of water on TiO2 prospers in last several decades, particularly
detailed understandings of water adsorption, diffusion, and dissociation on prototypical
TiO2 rutile(110).[49] Besides that, the study of the interaction of water with TiO2 surface
is inevitable to fully understand the chemical and catalytic properties of TiO2, because
water always be present either as part of an aqueous environment or as water vapor,
even in UHV conditions.
Many research demonstrate that dissociative adsorption of water happens at OV sites
on reduced TiO2 surface, followed by the appearance of OHP located on two
neighboring Obr sites, with the details of the dissociation of water shown in Figure
1.6.[50] The dissociative adsorption of water molecules and the tip assisted dissociation
of water molecules and OH have also been reported by many groups.[51-56] While on
oxide TiO2 surface, not only the dissociative adsorption of a water molecular is
observed at an OV site, but also the formation of a new water molecule at 300 K, as
revealed in Figure 1.7 by STM.[57]
Other TiO2 surfaces have also draw abundant attentions, for examples, rutile(011),
(100) and (001), as well as anatase(101) and (001), because the phase and facet are
recognized to be closely related to the (photo)catalytic performances.[58,59] It is also
known that the mixed TiO2 of rutile and anatase even exhibit much better
photocatalytic activity than the individual polymorphs, due to the proper band
alignment could effectively promote the separation of photoexcited charge
carriers.[60,61] Among these studies, anatase(101), which is the most frequently exposed
surface of this high active TiO2 polymorph, has been widely studied.[35, 62-68] As shown
in Figure 1.8, water monomer appears as isolated black spot surrounded by two bright
features (whit-black-white, w-b-w). In addition, through STM it was found that at 190
11
K, water monomer appearing as w-b-w can hop on the surface, as shown in Figure 1.8
(c)-(f). With more water on the surface (0.24 ML), they tend to form a (2 × 2)
superstructure on the surface, induced by the charge rearrangement at the molecule-
anatase interface.[62]
Figure 1.8 (a) Schematic diagram of the crystal structure of anatase(101) surface. (b)
STM image of anatase(101) surface at empty state, with unit cell marked by the black
rectangle containing two equivalent Ti5c/O2c surface atoms, and a water monomer
marked by the black arrow. (c) and (d) two consecutive STM images of anatase(101)
surface with 0.11 ML water on the surface, showing the hopping of water monomer
during STM scanning. (e) Difference image of c and d. (f) Height profiles crossing a
water molecule, and two adsorbed water molecules along the blue and red line in d,
respectively. (g) Ordered water overlayer in (2 × 2) on anatase(101), as indicate by the
red unit cell.[62] Reproduced with permission from Springer Nature.
1.2.3 Semiconductor photocatalyst and photo(electro)catalysts
Photocatalysis, through which one can convert solar energy into chemical energy,
has been regarded as one of the most promising strategies to solve crisis of energy
shortage and environmental pollution. It has been demonstrated to exhibit great
potential in several applications, including photocatalytic water splitting,
photosynthesis, and the treatment of pollutants in aqueous or gas phase.[69-72]
Unfortunately, reliable photocatalysts are scarce, although the appeal of the direct
12
conversion of solar into chemical energy via semiconductor compounds has been
recognized for a long time. This is mainly due to their low solar-energy conversion
efficiency, especially in the visible-light spectrum.
Generally, a typical photocatalytic process involves three steps: (i) absorption of
solar light (with the efficiency denoted as ηA), (ii) separation of photoexcited charge
carriers (with the efficiency denoted as ηS), and (iii) reduction or oxidation reactions at
the surface to complete the solar energy conversion (with the efficiency denoted as ηC).
The overall efficiency (η) is therefore determined by multiplying the efficiencies of the
individual steps: η = ηA × ηS × ηC. Thus, several criteria are essential for the design and
development of novel photocatalysts that possess high solar-energy conversion
efficiency. First of all, the band gaps of the photocatalysts should be narrow enough
(Eg < 3.0 eV), which allow absorption of both ultraviolet (UV) and visible light in the
solar spectrum. This is because UV and visible light account for 5% and 43% of the
solar spectrum, respectively. Secondly, a high separation rate and high mobility of the
photoexcited charge carriers (electrons and holes) are essential for high-efficiency
photocatalysts. The former can lead to high quantum conversion efficiency, while the
latter increases the effective charge-carrier diffusion length and thus enhances the
photocatalytic activity. For semiconductors, the mobility of electrons and holes is
determined by their effective mass (m*). As shown in Equation 1.1, the m* is estimated
to the second order derivative of energy (E) with respect to the wave vector (k), which
is reflected by the curvature of the band edges.
1
𝑚∗ =1
ħ2 × 𝑑2𝐸
𝑑𝑘2 (1.1)
As illustrated in Figure 1.9 (a), a larger curvature of the band leads to a small m*
(light effective mass) of the charge carriers. This means the dispersion of the electronic
band determines m*, that is, the more dispersive the band the smaller m* will be, and
13
consequently, the higher the mobility it will have. In contrast, a dispersion-less band
leads to large m* (heavy effective mass) of charge carriers with low mobility. Thirdly,
the positions of the valence and conduction bands (VB and CB) are of importance to
determine the chemical potential of photoexcited electrons and holes, which have
significant impact on the efficiency of different photocatalytic reactions (Figure 1.9
(b)). The overall solar-energy conversion during the photocatalysis process is,
therefore, determined by the electronic structures of photocatalysts.
Figure 1.9 (a) Diagram of the relationship between the band shape and the effective
mass of electrons and holes. (b) Diagram of the band structure of direct-gap
semiconductors, illustrating its role in determining the catalytic properties of
photocatalysts.
Surface and interface properties of semiconductors play a vital role in determining
their photocatalytic activity. This is not only because nanostructure or microstructure
with very large specific surface area are desired, but also due to electronic structure of
photocatalysts and separation of photo-excited electron-hole pairs can be effectively
tuned through surface and interface engineering.[73]
14
1.2.4 BiOBr based catalysts
BiOX (X = Cl, Br, I and F) belongs to the family of bismuth oxyhalides (BixOyXz).
As a heavy metal, bismuth atoms tend to form very weak bonds with non-metal
elements, which generally tend to have small molecular orbital overlap and favors the
establishment of narrow band gaps. Along with a narrow band gap, a small effective
mass may to be acquired for semiconductors with similar crystal and band structures,
according to the k·p perturbation theory. Various bismuth oxyhalides compounds
exhibit high visible-light activity towards photocatalytic degradation of organic and
inorganic toxic substances, water splitting, CO2 reduction and N2 fixation.[74-84]
Their VB maximum is mainly composed of O 2p and X np states (n = 3, 4, and 5 for
Cl, Br, and I, respectively). Their CB minimum in most cases is constructed from Bi
6p states.[85-87] Thus, the band structures of bismuth oxyhalides are expected to be
tunable via the halogen species and the ratios of Bi:O:X. Meanwhile, due to the
dispersive properties of the p and s-p hybridization states, high-mobility charge carriers
can be obtained in many bismuth oxyhalides compounds.
As shown in Figure 2.0 (a) and (b), both experimental and theoretical works have
suggested that the band gap of BiOX can be gradually narrowed from 3.4 eV for BiOCl
to 2.8 eV for BiOBr and 1.9 eV for BiOI, due to the increasing participation of the X
np states.[88,89] As a result, Zhang et al. proved that BiOI shows a wider visible-light
absorption range than BiOCl and BiOBr, and higher visible-light photocatalytic
activity towards degradation of methyl orange (MO).[90] On the other hand, due to the
appropriate VB edge, BiOBr and BiOCl exhibited much higher oxygen evolution
efficiency and better photocatalytic degradation of Rhodamine B (RhB) and phenol
under visible light or simulated sunlight irradiation than BiOI.[88,91] As shown in Figure
2.0 (d) and (e), Bhachu et al. demonstrated that BiOBr also had much superior water
15
oxidation activity compared to BiOI and BiOCl under simulated sunlight irradiation,
due to its more positive VB edge.[92] It should be noted that for BixOyXz photocatalysts,
the nature of their internal electric field contributes to the efficiency of separating
photoexcited charge carriers, which is induced by their layered structure, consisting of
interleaving positive [Bi-O] layers and negative X-layers.[93,94]
Figure 2.0 (a) Ultraviolet-visible (UV-Vis) diffuse reflectance spectra of BiOX (X =
Cl, Br, and I).[85] (b) Calculated band gap and the band alignment of the BiOX (X = Cl,
Br, I and F) compounds by density functional theory (DFT).[89] (c) Oxygen evolution
of BiOX under simulated sunlight irradiation.[29] (d) Photoanodic activity measurement
of BiOBr film under simulated sunlight irradiation. (e) Stability of the BiOBr film at
an applied voltage of 1.0 V vs. Ag/AgCl.[92]
Adjusting the ratio of Bi:O:X in bismuth oxyhalides has also been demonstrated to
be an effective approach to obtain visible-light photocatalysts with the desired
electronic structure.[76,77,95,96] Shang et al. reported that Bi24O31Br10 (Eg ≈ 2.8 eV,
16
similar to BiOBr) exhibited considerable photocatalytic activity towards Cr(VI) ion
reduction and H2 evolution through water splitting. BiOBr cannot split water to release
H2 through photocatalytic reactions, however, because its conduction band minimum
(CBM) is more positive than the electrode potential of H+/H2. While in the case of
Bi24O31Br10, the CBM, mainly consisting of hybridized Bi 6p and Br 4s orbitals, is
uplifted to be more negative than the electrode potential of H+/H2, which enables
Bi24O31Br10 to reduce water into H2 under visible-light irradiation. In addition, the
hybridization s-p orbitals lead to a dispersive band, which is expected to increase its
efficiency through promoting the separation of the photoexcited electrons and holes.
Surface engineering has successfully been implemented in promoting the
photocatalytic activity of bismuth oxyhalides. Manipulations of exposed facets have
been achieved in a variety of BiOX photocatalysts, in which the tunable exposed facets
can efficiently adjust their photocatalytic activities.[75,97-101] Morphology and thickness
engineering are also efficient approaches for modulating the photocatalytic activities
of bismuth oxyhalides through adjusting their the exposed surfaces, specific surface
area and electronic properties, which are closely relevant to their layered structures.[102-
104]
In addition, OVs have been reported to be particularly efficient in promoting the
photocatalytic activity of ultra-thin bismuth oxyhalides nanosheets. For examples,
enhanced photocatalytic performances in dye degradation by BiOCl,[105-107] and CO2
reduction by BiOBr,[74, 108] through introducing OVs in their ultra-thin nanosheets. It
demonstrates that effect of defect engineering could be more effective for
semiconductor photocatalysts in 2D, which has a very high specific surface area and
more sensitive electronic properties compared with bulk materials.
17
1.3 Outline of chapters
Chapter 2 introduces the experimental instruments used in this thesis. SPM
techniques, comprising AFM and STM, are illustrated in details of their working
principles and applications. Other facilities, including electron microscopies (SEM,
TEM and STEM) and photoelectron spectroscopy (XPS and UPS), which are widely
applied in the study of surface science and technology, are also introduced.
Chapter 3 presents STM study of the evolution of the TiO2 rutile(110) surface with
increasing oxygen deficiency, which was created by annealing in vacuum condition.
At 900 K, the surface exhibit a (1×1) structure with isolated OV point defects. While
at 1300 K, the surface changed to (1×2) reconstruction with cross-links defects
dominating the surface.
Chapter 4 presents the surface study of OV-TiO2(110) surfaces by STM and STEM.
The effect of OVs on the electronic properties were further revealed by photoelectron
spectroscopy. More importantly, the vital roles of OVs on the greatly enhanced
electrocatalytic activity toward HER were revealed under the combination of surface
studies and DFT calculations. The promoted electrical conductivity and hydrogen
desorption capability induced by OVs were suggested to be decisive in the enhanced
electrocatalytic activities.
Chapter 5 presents the study of modification of inner strain in 2D {001} facets
exposed BiOBr nanosheets. The effect of strain on the crystal structure and electronic
structure were revealed by TEM and DFT calculations. It is also demonstrated that the
photocatalytic performance of BiOBr in dye degradation can be manipulated by the
strain effect. The improved charge separation attributable to a highly dispersive band
structure in low-strain BiOBr nanosheets contributed to the promoted photocatalytic
activities.
18
Chapter 6 presents the strain induced 2D lateral pseudoheterogeneous structures in
2D BiOBr nanosheets. AFM, TEM and STEM characterizations revealed that the
pseudoheterogeneous interface without atomic mismatch can be feasibly modulated by
local strain distribution, which exhibits similar local electronic band structure of
corresponding heterostructures. Significant enhancement in charge separation at the
pseudoheterostructure was demonstrated under visible light irradiation of individual
single nanoplates, which is given rise to the controllable electronic band alignment
across the interface. The construction of the lateral pseudoheterostructure offers a
feasible and promising way to build unprecedented 2D systems with exciting
properties.
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32
Chapter 2 Experimental Techniques
2.1 Scanning tunneling microscopy
STM was firstly demonstrated in 1982 by Gerd Binning and Heinrich Rohrer, who
were later awarded the Nobel Prize in physics in1986 for their design of STM.[1,2] STM
can achieve atomic resolution on the conductive surface of semiconductors and metals,
and is able to acquire the local density of states (LDOS) as function of position on the
surface. In addition, STM has also demonstrate its strong capability of manipulating
and transferring atoms on the surface. As a consequence, STM is applied in a wide
range of research and application fields, including physical, chemical, biological,
material, and nano science and technology.
2.1.1 Quantum tunneling
The quantum tunneling effect is a quantum phenomenon that has no counterpart in
classical physics. A particle has a probability to tunnel cross a barrier, even though the
energy of the particle is less than the barrier height. These barriers can either be
physically impassable medium, including insulators or vacuum, or a region of high
potential energy.
33
Figure 2.1 Quantum mechanical tunneling through a potential barrier of height (V) and
width (D).
The quantum tunneling probability of a particle can be determined through solving
Schrödinger equation. As shown in Figure 2.1, the incident particle from the left of the
energy barrier (width D) has a free particle wavefunction of an energy smaller than the
potential barrier of height (V). When the incident particle approaches the barrier,
The Schrodinger equation can be described as:
−ħ2
2𝑚 𝜕2Ψ(x)
𝜕x2 = (𝐸 − 𝑉)Ψ(x) (2.1)
The solution is:
Ψ = A𝑒−𝛼𝑥 (2.2)
where the decay constant α: 𝛼 = √2𝑚 (E−V)
ħ2 (2.3)
Therefore, the decay constant depends not only on the energy barrier but also the
mass (m) of the particles and the reduced Planck constant (ħ).
34
2.1.2 Working principle of STM
STM is designed based on the principle of quantum tunneling of electrons between
a sharp metallic tip and the conductive surface, when they have a close distance within
several angstroms. Meanwhile, in order to achieve the scanning of the surface, a 3-
dimisonial piezoelectric scanner with the sub-angstrom precision in both x-y plane and
z direction are need. When a voltage between the tip and the sample is given, electrons
can tunnel between them. For example, if a positive voltage is applied on the sample,
electrons will tunnel from the filled state of the tip to the empty states of the sample
when they are in the tunneling region, and vice versa. The tunneling current can be
calculated using the time-dependent perturbation theory. If a positive V is applied to
the sample, the Fermi level of the sample shifts down with respect to the Fermi level
of the tip, and electrons tunnel from the occupied states of the tip into the empty states
of the sample. The net electric current, which is in the range of picoamperes to
nanoamperes, then can be detected. The intensity of the tunneling current is
proportional to the overlap between the wavefunctions of tip and sample, which
depends strongly on their distance. A typical setup of STM system is illustrated in
Figure 2.2.
Figure 2.2 Diagram of the setup of STM system.
35
The total tunneling current can be estimated by the perturbation theory[3-5]:
𝐼(𝑉) ∝ ∫ 𝜌𝑠 (𝐸)𝜌𝑡(𝐸 − 𝑒𝑉)|𝑀𝑡,𝑠(𝐸, 𝑉, 𝑧)|2
(𝑓(𝐸 − 𝑒𝑉, 𝑇) − 𝑓(𝐸, 𝑇))𝑑𝐸 (2.4)
where ρs is the DOS of sample, ρt is the DOS of tip, E is the energy of electrons, T
is temperature, f is Fermi function, V is the applied voltage, 𝑀𝑡,𝑠(𝐸, 𝑉, 𝑧) is the tunnel
matrix.
According to Tersoff-Haman approach, ρs and ρt can be treated as constant under
small bias voltage. Considering in STM studies, low temperature and metallic tips
(such as W or PtIr) with flat DOS around Fermi level are applied, the expression for
the tunneling current can be simplified as:
𝐼(𝑉) ∝ ∫ 𝜌𝑠 (𝐸)𝑑𝐸 (2.5)
Therefore, the tunneling current under certain bias V is proportional to the integral
of the DOS of the sample from the Fermi level to eV.
For the widely adopted constant-current mode, a feedback loop is used to regulate
the distance between STM tip and the sample in the z-direction to maintain the set-
point current. Then the change of the z position of the tip can be obtained to produce
the surface topography.
Meanwhile, tunneling current is dependent on distance between STM tip and sample,
and the integral of the DOS from Fermi level to eV. For a sample with a homogenous
DOS, the topography is entirely contributed by the geometric surface profiles.
Whereas, considering that most materials exhibit a spatially inhomogeneous DOS, the
obtained STM images are dominated by both the geometric surface profile and the local
DOS.
As the measured tunneling current is proportional to the integral of the DOS of the
sample from the Fermi level to eV, STM can also acquire the local DOS (LDOS) of a
36
selected energy range on the surface by fixing distance between STM tip and the
sample. This means the LDOS can be directly measured by sweeping the applied
voltage at a fixed position on the surface, which can be explained by Equation 1.6
(Derived from Equation 2.5). Instead of numerical calculation of the obtained tunneling
current and the applied voltage that is easily be affected by the noise, lock-in amplifier
technique is widely employed to directly record the dI/dV signal, which works through
applying a small bias voltage modulation dV to the applied V and record the change in
the tunneling current dI.
𝑑𝐼
𝑑𝑉∝ 𝜌𝑠(𝑒𝑉) (2.6)
2.1.3 Ultra-high vacuum system and molecular beam epitaxy
In the LT-STM and MBE combination system, UHV condition is necessary for
acquiring highly clean sample surface, and high resolution and stable images. To
achieve UHV (less than 1× 10-10 Torr) in our system, a serious of supporting structures
and vacuum pumps are used, as illustrated in the setup diagram in Figure 2.3.
Figure 2.3 The relative schematic vacuum layout of our STM/MBE system.
A load-lock is designed for loading and transferring samples and STM tips into the
UHV preparation chamber without venting the chamber to atmosphere. The vacuum of
37
load-lock chamber is achieved by to 10-9 Torr. Load-lock chamber is separated from
preparation chamber and observation chamber by a gate value (GV). To obtain the a
UHV condition for preparation chamber and observation chamber, the system needs to
be firstly pumped by a combination of rotary pump and turbo molecular pumps (TMP)
to the vacuum level around 1 × 10-9 Torr after baking (generally 100 oC to 250 oC).
Then under the pumping of ion pump and titanium sublimation pump (TSP), the system
can reach an UHV condition of 5 × 10-11 Torr.
Figure 2.4 The relative schematic of preparation chamber with MBE system.
Molecular beam epitaxy (MBE) is an advanced UHV epitaxy method with high
precision of sub monolayer and purity. The key aspect of MBE is the slow deposition
rate, which enable the precisely control of the amount of dopants or thickness of crystal
film. In the UHV environment, atom or molecular beams are thermally evaporated and
deposit on a heated substrate placed in the line of sight of the beam line, as shown in
38
the relative schematic of our preparation chamber. The widely adopted evaporation
source in commercial MBE is Knudsen Effusion Cell (K-cells), which consists
mechanical shutters, boron nitride or aluminum oxide crucible, tungsten filament
heater, tantalum heat shielding, thermocouple, and cooling water system. Other
facilities, for examples the heating stage and the argon sputtering gun, are also installed
in the preparation chamber for both sample deposition, treatment and cleaning.
2.2 Atomic force microscopy
Atomic force microscopy (AFM), which is another important SPM technique
beyond STM, was developed by Gerd Binning et al. in 1986.[6,7] Different with STM
which was designed based on the quantum tunneling effect between STM tip and
sample surface, AFM works through measuring the force interaction between AFM tip
and sample surface. Thus, the application scopes of AFM are not limited for conductive
samples, but also can be extended to almost any measurable force interaction, including
van der Waals, electrical, chemical bonding, magnetic, thermal, and capillary forces,
which enable AFM to be the most versatile SPM technique.[8] Thus, combining the
high-resolution for 3D topography, capability of operating in ambient conditions and
liquid conditions, and integrating with variety of optical microscopy and spectroscopy
techniques, AFM has been widely employed in measuring surface topography and
detecting force interaction in many research fields, such as material science, condensed
matter physics, nanoscience and nanotechnology, chemistry, biology, and medicine.
2.2.1 Working principles of AFM
AFM, typically, mainly consists of a sharp tip fixed on a cantilever, laser beam and
optical system, laser detector and feedback system, display and processing system, and
piezo-based scanning stage, as shown in Figure 2.5.
39
Figure 2.5 Diagram of the setup of a basic AFM in laser beam deflection system.
𝐹𝑡𝑠 = −𝜕𝑉𝑡𝑠
𝜕𝑧 (2.7)
where Fts is the tip-sample force and Vts is the potential.
When a sharp AFM tip approaches close a sample surface, a force will arise between
AFM tip and the sample due to the potential energy, which is illustrated in Equation
2.7. Then the force lead to a deflection of the cantilever according to Hooke's law, with
the spring constant kts determined as:
𝑘𝑡𝑠 = −𝜕𝐹𝑡𝑠
𝜕𝑧 (2.8)
Fts can be divided into long-range and short-range forces. For the ambient AFM
system employed in this thesis, long-range force typically represents the attractive
forces, including Van der Waals force, magnetic force, capillary force, and electrostatic
force. While, short-range force can be measured in vacuum conditions when the tip-
sample distance can reach less than several angstroms), chemical force, covalent force,
electrostatic, magnetic and Van der Waals forces can be determined.
2.2.2 Working modes of AFM
There are basically three operation modes of AFM classified by the way of tip
motion, including contact mode, tapping mode and non-contact mode. These different
40
working modes can be adopted in different interaction forces, and corresponding
applications. Figure 2.6 illustrates the relationships between AFM working modes and
the force ranges. Normally, contact mode and tapping mode are used in ambient
environment, while non-contact mode is mostly used in vacuum. Therefore, only
contact mode and tapping mode are introduced in this thesis.
Figure 2.6 Classification of AFM working modes based on Van der Walls force curve.
Contact mode-AFM
In contact mode, the AFM tip directly contact with the sample surface under
repulsive force. The constant force mode is mostly adopted for most applications.
During the scanning, the deflection of the cantilever is kept constant, which enables the
force to be constant. The constant deflection of the cantilever is controlled by an active
feedback loop that can adjust the tip and sample surface distance corresponding to the
topography. For constant height mode, the spatial variation of the cantilever deflection,
in contrast, is directly recorded to yield the topography of sample surface. Constant
41
height mode is mostly employed only for acquiring atomic resolution of sample
surface. In addition, based on contact mode AFM, a serious of functions have been
developed, for examples, force spectroscopy and mapping, lithography and
nanomanipulation, conductive AFM, and so on.
Tapping mode-AFM
In tapping mode, the AFM tip oscillates above sample surface under a driving
frequency very close to the resonance frequency of the cantilever. During the scanning,
the oscillation amplitude of the cantilever is kept constant through the feedback loop,
when the tip periodically contacts with sample surface. As the force between the tip
and sample surface changes, the feedback loop will adjust the height to maintain the
set oscillation amplitude of the cantilever, which reflects the force distribution or the
topography of the sample surface. In tapping mode, as the tip gently interact with
sample surface, both the tip and sample are less destructive compared with contact
mode. In addition, in tapping mode, the phase lag of the cantilever oscillation relative
to the signal sent to the driving piezo can be simultaneously recorded. The phase
channel reflects the energy dissipated by the cantilever during the scanning, which can
reveal the stiffness or adhesion information of sample surface.
Kelvin probe force microscopy (KPFM) combined with Tapping mode
KPFM, also known as surface potential imaging, has been well developed combined
with tapping mode AFM, in which the difference between the potential of the tip and
that of the sample can be determined.[9,10] The data obtained in this mode is a
combination of three contributing factors: the work function difference, trapped charge,
and any permanent or applied voltage between the tip and the sample. As a consequent,
42
KPFM is generally considered a pseudo-quantitative technique, in which the acquired
accurate contact potential difference (CPD) mostly consists of more than one physical
quantity. In KPFM measurement, an AC bias (VAC) is usually applied on a conductive
tip to oscillate the electrostatic force between tip and sample. The feedback loop then
can provide a DC bias (VDC) to the tip to minimize the electrostatic force between the
tip and the sample.
The electrostatic force between the tip and the sample under the VAC can be described
as Equation 2.9, when they are modeled as a parallel plate capacitor.
𝐹 =1
2
𝜕𝐶
𝜕𝑧𝑉2 (2.9)
The total potential difference (V) between the tip and the sample is the sum of VAC,
VCPD, and the VDC, as shown in Equation 2.10.
𝑉 = 𝑉𝐶𝑃𝐷 + 𝑉𝐷𝐶 + 𝑉𝐴𝐶𝑠𝑖𝑛(⍵𝑡) (2.10)
Therefore, the electrostatic force can be described as:
𝐹 =1
2
𝜕𝐶
𝜕𝑧([(𝑉𝐷𝐶 − 𝑉𝐶𝑃𝐷)2 +
1
2𝑉𝐴𝐶
2 ] + 2[(𝑉𝐷𝐶 − 𝑉𝐶𝑃𝐷)𝑉𝐴𝐶sin (⍵𝑡)] −
[1
2𝑉𝐴𝐶
2 cos (2⍵𝑡)]) (2.11)
The first part represents force is static and not frequency dependent. The second part
occurs at the drive frequency ⍵. The third part reacts at twice the drive frequency. As
a result, the most important term here as far as surface potential is concerned is the
second, since this depends not on the square of the voltage, but rather on the potential
difference between the tip and the sample, multiplied by the magnitude of the applied
VAC. This means that if there is a VCPD between the tip and the sample, then when an
VAC voltage is applied, there will be an oscillatory force at the frequency of the drive
and proportional of the magnitude of the applied voltage, and also proportional to the
CPD. Further, if we make VDC = VCPD, with a potential feedback loop between the tip
43
and the sample, then the oscillations at ⍵ will be nulled. Therefore, the CPD can be
recorded during the scanning.
2.3 Other facilities
Besides STM and AFM, which can directly probe the surface morphology, and in
some extent, visit the surface electronic states and catalytic reactions, other facilities
have also demonstrated their capabilities in this research filed. Here, several facilities
of electron microscopies, X-ray and electron spectroscopies involved in this work are
briefly introduced.
2.3.1 Scanning electron microscopy and focused ion beam
Scanning electron microscopy (SEM) is a very useful and robust facility for imaging
the surface morphology of samples with a resolution higher than 1 nanometer, which
belongs to the family of the electron microscopes. It works through scanning the
electrically conductive sample surface with a focused electron beam. For samples that
are not conducting, coating with very thin noble metal film (Au or Pt) enable their
surface morphologies to be imaged by SEM. During the interactions between the
focused electron beam and sample surface, main products are secondary electrons,
backscattered electrons, auger electrons, and characteristic x-rays. These products can
be collected by different detectors, which contribute the SEM images. For examples,
secondary electrons come from the inelastic interactions between the primary electron
beam and the sample, which mostly contributed by the atoms in top surface and near
surface regions. It is the most common used SEM mode in imaging the surface
morphology. The characteristic x-rays signal can be collected for the energy-dispersive
x-ray spectroscopy (EDX), which enables elemental analysis or chemical
characterization of sample surface.
44
Besides SEM, focused ion beam (FIB) is another important technique in the fields
of materials science and industry. Different from SEM which is designed based on
focused electron beam, FIB commonly uses a focused ion beam, or integrate focused
electron beam together. Therefore, FIB is widely used for lithography and etching in
micro or nano-scale by the focused high energy ion beam (Gallium ion is most widely
used ion source). In addition, FIB has also been used in locally depositing films,
modifying and fabricating materials, and preparing transmission electron microscopy
(TEM) specimen.
2.3.2 Transmission electron microscopy
TEM is another very important electron microscopy, which is well known as its
capability of imaging real-space atomic-resolution of nanostructures. TEM uses high
energy focused electron beam to transmit through and interact with the specimen as it
passes through it to image the structure, shape, morphology, size and composition of
the specimen. The specimen used in TEM are required to be have a thickness less than
100 nm or in nanostructures that can be suspension on a grid, which guarantee effective
transmission of electron beam. TEM appears in three different forms of high-resolution
transmission electron microscopy (HRTEM), scanning transmission electron
microscopy (STEM) and analytical electron microscope (AEM).[11]
The working principle of TEM can be described using a single lens microscope, with
the difference in using electrons rather than photons as the source. Figure 2.7 shows
the simplified model, in which only the objective lens that determining the resolution
are included and the intermediate lenses and projection lenses are omitted. The
resolution of TEM can be estimated by the Abbe’s equation modified by the by using
DeBroglie’s formula.
45
𝑑 =0.753
𝛼𝑉12
(2.12)
where d is the resolution in nanometer, α is the half aperture angle, V is the
accelerating velocity of the electron beam. For a TEM system with accelerating voltage
of 100 kV, the resolution d is valued to be 0.24 nm.
Figure 2.7 The simplified model of a one-lens TEM based on the Abbe’s theory.[12]
HRTEM is one of the most well-developed tools to in material science with its
atomic resolution capability. HRTEM uses both the scattered and transmitted electrons,
which is an interference pattern between the forward-scattered and diffracted electron
46
waves from the specimen. Therefore, it is a phase contrast imaging mode, which can
achieve a resolution as high as 0.1 nm. The acquired diffraction pattern with the
information of the specimen, then can be converted to the Fourier transform analysis.
In recent years, STEM has experienced revolutionary development due to the
aberration correction technique. Sub-angstrom resolution has been achieved in
STEM,[13] which greatly benefits many research filed, such as the solid state physics,
materials sciences, catalysis and chemistry sciences. In addition, the successfully
implementation of combining high angle annular dark-field (HAADF) mode and
electron energy loss spectroscopy (EELS) which belongs to AEM have been achieved.
This provides a very powerful approach to coupling high resolution images with the
local electronic structure of the specimen, which are valuable for revealing fundamental
mechanism for many scientific challenges, especially in the field involving surface and
interface.
Different with HRTEM, STEM uses the focused electron to scan over the specimen.
The electron beam is focused with the assistance of the aberration correction to reach
a fine spot in the order of sub-angstrom. The scattered electrons undergone different
interactions are then collected by the different detectors. The most common detectors
in commercial STEM include HAADF, annular dark-field (ADF), bright-field (BF),
EELS and EDX. The HAADF detector is, in particular, important in acquiring atomic
resolution images. It collects the electrons scattered out to high angles, in which
electrons are not Bragg scattered. As a result, HAADF images show little or no
diffraction effects. The detected intensity of the incoherently scattered electron is
highly sensitive with the atomic number and the image contrast has a proportional
relationship Z2 according to Rutherford scattering model. Therefore, the extremely high
47
atomic resolution is determined by the size of the focused electron spot, rather than
HRTEM which depended on electron diffractions.
2.3.3 Photoelectron spectroscopy
Photoelectron spectroscopy is the most widely used quantitative spectroscopic
technique in analyzing the surface chemical and elemental state, and electronic
structure of materials. Photoelectron spectroscopy is designed based on the
photoelectric effect, using photo-ionization and measuring the kinetic energy
distribution of the emitted photoelectrons. The recorded photoelectron spectrum then
can be converted to the binding energy of electrons which represents the energy
difference between the ionized and neutral atoms. The relationship is described by the
Einstein relationship:
𝐸𝑏 = ℎν − 𝐸𝑘 − (2.13)
where Eb is the electron binding energy, hν is the photon energy of the radiation
source, Ek is the kinetic energy of photoelectron, is the work function induced by the
analyzer, which dependent on both the analyzer and the material. This equation is
essentially a conservation of energy equation. The work function, , usually can be
regarded as constant in the measurement.
According to the different exciting radiation source, photoelectron spectroscopy can
be divided into x-ray photoelectron spectroscopy (XPS) and ultraviolet photoelectron
spectroscopy (UPS). As shown in the diagram in Figure 2.8, XPS normally uses soft
x-rays (200-2000 eV) as the radiation source and can examine the core-level electrons
with a penetration depth of 1-10 nm. While UPS applies vacuum UV radiation (10-45
eV) to examine valence elections of samples with a penetration depth even smaller than
XPS due to its lower incident photon energies.
48
Figure 2.8 Schematic diagram of the electron excitation progress of XPS and UPS.
It has to be mentioned that, the implementation of synchrotron radiation as the
excitation source for photoelectron spectroscopy has greatly broaden its
applications.[14,15] Compared with the normal x-ray generator, and He lamp for UV
light, synchrotron radiation has advantages of high intensity, very broad and continuous
spectral range, high flux and brightness, and high degree of polarization, which has
been playing a very important role in many advanced research fields, including
materials science, physical, biological and chemical sciences, geosciences,
environmental sciences, and medical and pharmaceutical sciences.
49
2.4 References
(1) G. Binning, H. Rohrer, C. Gerber, E. Weibel, Surface studies by scanning tunneling
microscopy, Phys. Rev. Lett. 1982, 49, 57.
(2) G. Binning, H. Rohrer, Scanning tunneling microscopy from birth to adolescence,
Rev. Mod. Phys. 1987, 59, 615.
(3) C. J. Chen, Introduction to Scanning tunneling microscopy, Oxford Univ. Press,
New York, 1993.
(4) J. Tersoff, D. R. Hamann, Theory and application for the scanning tunneling
microscope, Phys. Rev. Lett. 1983, 50, 1998.
(5) J. Tersoff, D. R. Hamann, Theory of the scanning tunneling microscope, Phys. Rev.
B 1985, 31, 805.
(6) G. Binnig, Atomic force microscope and method for imaging surfaces with atomic
resolution, 1986, US Patent No. 4, 724, 318.
(7) G. Binning, C. F. Quate, C. Gerber, Atomic force microscope, Phys. Rev. Lett. 1986,
56, 930.
(8) F. J. Giessibl, Advances in atomic force microscopy, Rev. Mod. Phys. 2003, 75,
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(9) M. Nonnenmacher, M. P. O’Boyle, H. K. Wickramasinghe, Kelvin probe force
microscopy, Appl. Phys. Lett. 1991, 58, 2921.
(10) V. Palermo, M. Palma, P. Samorì, Electronic characterization of organic thin films
by kelvin probe force microscopy, Adv. Mater. 2006, 18, 145.
(11) D. B. Williams, C. B. Carter, Transmission electron microscopy, Plenum press,
New York, 1996.
(12) Z. L. Wang, Transmission electron microscopy of shape-controlled nanocrystals
and their assemblies, J. Phys. Chem. B 2000, 104, 1153.
50
(13) O. L. Krivanek, M. F. Chisholm, V. Nicolosi, T. J. Pennycook, G. J. Corbin,
N. Dellby, M. F. Murfitt, C. S. Own, Z. S. Szilagyi, M. P. Oxley, S. T. Pantelides, S. J.
Pennycook, Atom-by-atom structural and chemical analysis by annular dark-field
electron microscopy, Nature 2010, 464, 571.
(14) C. S. Fadley, X-ray photoelectron spectroscopy: progress and perspectives, J
Electron Spectrosc. 2010, 178-179, 2.
(15) S. Günther, B. Kaulich, L. Gregoratti, M. Kiskinov, Photoelectron microscopy and
applications in surface and materials science, Prog. Surf. Sci. 2002, 70, 187.
51
Chapter 3 Characterization of TiO2(110) surface by STM
3.1 Introduction
TiO2, as a typical TMO, have been widely applied in the field of energy conversion,
surface pigments, and photocatalytic wastewater and hazardous gas remediation, due
to its catalytic activity and semiconducting properties as well as its earth abundance,
low toxicity, chemical and thermal stability.[1-2] Surface and interface engineering has
been recognized as an effective approach in promoting the performances or tuning
reactivity of TiO2 in catalysis or energy conversion, for examples facet engineering,[3,4]
curved surfaces,[5] surface defects,[6-8] surface adsorbates and constructing interface
heterostructures[9-14].
Among these strategies, OV is regarded as one of the most important and prevalent
defects in TiO2, as well as a large number of metal oxides, which has been widely
investigated both by theoretical calculations and experimental characterizations.[15-20]
Compared with defects created by doping, OVs are the kinds of defect which do not
disturb the intrinsic crystal structure and involve other impurity elements. At the same
time, OVs are thought could contribute excess electrons that will occupy 3d orbitals of
the neighboring Ti atoms and forms Ti3+ ion, which could enhance the electronic
conductivity of TiO2. As a result, OVs are expected to significantly affect the physical,
chemical and catalytic properties of TiO2, for examples, absorbing and desorbing
behaviors toward molecules, electron transfer in the catalyst and between the surface
of catalyst and adsorbed molecules, as well as the ability of light absorption.
Therefore, acquiring the detailed knowledges of OVs about their location,
distribution, as well as effect on electronic structure and adsorption/desorption of
52
molecules, especially on the surface region, are of great importance toward rationally
promote the activity of TiO2 and other TMOs in their applications in catalysis and
energy conversions. In this chapter, thermally created OV based defects were
introduced into (110) single crystal in vacuum conditions. By carefully controlling the
annealing temperature and time, the defect species and concentration on the surface
were examined by in-situ STM observations. It was found that at the surface of
rutile(110) experienced a reconstruction evolution from (1×1) with point OV defect to
(1×2) with cross-links defects, with the rutile(110) single crystal heated from 900 K to
1300 K.
3.2 Experimental section
TiO2 rutile(110) single crystal (5 × 5 × 1 mm) was purchased from Mateck, GmbH.
To obtain the reduced TiO2, the rutile single crystals were annealed in UHV at 900 K
to 1300 K. The concentration of OVs in the single crystal were controlled by the
annealing time and annealing temperature. STM images were acquired by using a low-
temperature STM (LT-STM) (USM 1500-M, Unisoku Co.) at 78 K. STM images can
be obtained in both constant current mode and constant height mode. Here, STM
images were acquired in constant current mode unless noted otherwise. The samples in
STM measurements were treated by several cycles of Ar+ sputtering (1 kV, 20 min)
and annealing.
53
3.3 Results and discussion
3.3.1 Rutile(110) surface with (1×1) reconstruction
Figure 3.1 Large scale STM image of the reduced TiO2 surface (1.3 V, 30 pA).
Figure 3.1 shows the large scale STM image of the rutile(110) surface after several
sputtering and annealing at 900 K, in which flat terraces were observed. When we
zoomed in to the very small scale shown in the empty-state images in Figure 3.2, atomic
resolution STM of the surface were acquired in both constant current mode and
constant height mode. As the empty-state of TiO2 is dominated by the Ti-3d electronic
states, 5f Ti atoms and Obr atoms appear as bright and dark rows in the empty-state
STM image, respectively, which is reverse-contrast of their real topographies. The
54
lattice constant of was measured to be 0.63 nm by 0.28 nm, with each bright spot
representing one 5f Ti atom, which agrees well with the crystal structure in Figure 1.3.
Figure 3.2 Atomic resolution of the stoichiometric rutile(110) surface. (a) constant
current mode, 0.5 V, 100 pA, scan size is 6 nm × 6 nm, (b) constant height mode, 0.5
V, 100 pA, scan size is 6 nm × 6 nm.
3.3.2 OV point defects on (1×1) surface
Figure 3.3 The reduced rutile(110) surface with OVs on the surface, 1.2 V, 20 pA.
55
In the STM image of reduced (sputtered and annealed at 900 K) rutile(110) surface,
bright protrusions between two 5f Ti rows can be observed, as indicated by the arrows
in Figure 3.3. The OVs trends to individually locate between 5f Ti rows. While as the
surface stored in the UHV chamber for several hours, other two kinds of bright
protrusions appeared on the surface, as shown in Figure 3.4 (a). These three types of
bright protrusions (marked as I, II, and III) exhibit different brightness and apparent
heights (indicated in Figure 3.4 (b)), which can be used not only in identifying different
species, but also in monitoring their dynamic reaction processes.[21-23]
Figure 3.4 (a) STM image of the surface of rutile(110) obtained several hours after
annealed in the UHV (10 nm × 10 nm, 1.6 V, 30 pA). (b) Apparent height profiles of
OV, OH group and H2O in STM images. OV, OH, and H2O are bright protrusions
between two adjacent Ti rows.
They can also be identified by their different responses toward voltage bias plus,
which a voltage higher than a thresh-old voltage (around 2.0 V in our measurements),
either through giving pulses or scanning, was applied. As shown in Figure 3.5, when
three 2.5 V pulses were applied at the selected OHs marked in Figure 3.5 (a), the
hydrogen atoms were removed, allowing the OV to be healed. In the case of OVs, they
are always inactive towards 2.5 V pulses or scanning. Therefore, by a combination of
their apparent heights and their different behavior under a bias higher the threshold
56
voltage, we can identify them. This phenomenon indicates that OVs on the surface of
rutile(110) are reactive toward water molecules even in UHV conditions and favorable
to be hydroxylated on the surface.
Figure 3.5 (a) STM image of the reduced rutile(110) surface with OVs and OHs, 1.3
V, 30 pA. (b) STM image in a same area in (a), in which 2.5 V pulses were applied at
the marked OHs, 1.3 V, 30 pA.
3.3.3 Rutile(110) surface with (1 × 2) reconstruction
Figure 3.6 Large scale STM image of the reduced TiO2 surface acquired after annealed
at 1300 K, 3 V, 50 pA, with the inset showing the side view of (1 × 2) reconstruction
surface.
57
When the annealed temperature was increased to 1300 K, the flat terraces on (1 × 1)
surface evolved to chain structures along Ti rows, as presented in Figure 3.6. This
surface has been explained by the (1 × 2) reconstruction due to the more oxygen
deficiency at higher annealing temperature.[24-26] In addition some short chains
perpendicular to the (1 × 2) chains can be observed, which are marked by the white
solid line as links in Figure 3.6. As shown in the zoomed-in images in Figure 3.7, the
unit cell of (1 × 2) reconstruction is marked by the black rectangle. Two cross-links of
single-link and double-link structure are indicated by the yellow and white arrows.
Meanwhile, it notices that the high quality STM image in filled state can be obtained
on (1 × 2) surface, which is not applicable on (1 × 1) surface, as shown in Figure 3.7
(b). It indicates that the electronic structures of rutile(110) also undergo significant
change with the increasing oxygen deficiency.
Figure 3.7 STM images of rutile(1 × 2) surface. (a) empty state mode, scan size is 10
nm × 10 nm, 1 V, 50 pA, (b) filled state mode, scan size is 10 nm× 10 nm, -1.8 V, 50
pA.
58
3.4 Summary
OV based defects were introduced into (110) single crystal through annealling in
vacuum conditions. By carefully controlling the annealing temperature and time, the
defect species and concentration on the surface were examined by in-situ STM
observations. On the rutile(110) (1×1) surfaces obtained at 900 K, isolated OV point
defects were found, which could be occupied by water molecule or convert to OH. On
the (1×2) surfaces obtained at 1300 K, cross-links defects dominated the surfaces.
3.5 References
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modifications, and applications, Chem. Sci. 2007, 107, 2891.
(2) B. O’Regan, M. Grätzel, A low-cost, high-efficiency solar cell based on dye-
sensitized colloidal TiO2 film, Nature 1991, 353, 737.
(3) H. G. Yang, C. H. Sun, S. Z. Qiao, J. Zhou, G. Liu, S. C. Smith, H. M. Cheng, G.
Q. Lu, Anatase TiO2 single crystals with a large percentage of reactive facets, Nature
2008, 453, 638.
(4) J. S. Chen, Y. L. Tan, C. M. Li, Y. L. Cheah, D. Luan, S. Madhavi, F. Yin, C. Boey,
L. A. Archer, X. W. Lou, Constructing hierarchical spheres from large ultrathin anatase
TiO2 nanosheets with nearly 100% exposed (001) facets for fast reversible lithium
storage, J. Am. Chem. Soc. 2010, 132, 6124.
(5) K. Shirai, G. Fazio, T. Sugimoto, D. Selli, L. Ferraro, K. Watanabe, M. Ηaruta, B.
Ohtani, H. Kurata, C. D. Valentin, Y. Matsumoto, Water-assisted hole trapping at the
highly curved surface of nanoTiO2 photocatalyst, J. Am. Chem. Soc. 2018, 140, 1415.
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(6) M. Kong, Y. Li, X. Chen, T. Tian, P. Fang, F. Zheng, X. Zhao, Tuning the relative
concentration ratio of bulk defects to surface defects in TiO2 nanocrystals leads to high
photocatalytic efficiency, J. Am. Chem. Soc. 2011, 133, 16414.
(7) M. Batzill, E. H. Morales, U. Diebold, Influence of nitrogen doping on the defect
formation and surface properties of TiO2 rutile and anatase, Phys. Rev. Lett. 2006, 96,
026103.
(8) X. Yu, B. Kim, Y. K. Kim, Highly enhanced photoactivity of anatase TiO2
nanocrystals by controlled hydrogenation-induced surface defects, ACS Catal. 2013, 3,
2479.
(9) G. Liu, L. Wang, H. G. Yang, H.-M. Cheng, G. Q. Lu, Titania-based photocatalysts-
crystal growth, doping, and heterostructuring, J. Mater. Chem. 2010, 20, 831.
(10) C. L. Pang, R. Lindsay, G. Thornton, Structure of clean and adsorbate-covered
single-crystal rutile TiO2 Surfaces, Chem. Rev. 2013, 113, 3887.
(11) J. Zhang, J. H. Bang, C. Tang, P. V. Kamat, Tailored TiO2−SrTiO3 heterostructure
nanotube arrays for improved photoelectrochemical performance, ACS Nano 2010, 4,
387.
(12) J. Tian, P. Hao, N. Wei, H. Cui, H. Liu, 3D Bi2MoO6 nanosheet/TiO2 nanobelt
heterostructure: enhanced photocatalytic activities and photoelectochemistry
performance, ACS Catal. 2015, 5, 4530.
(13) J. Resasco, H. Zhang, N. Kornienko, N. Becknell, H. Lee, J. Guo, A. L. Briseno,
P. Yang, TiO2/BiVO4 nanowire heterostructure photoanodes based on type II band
alignment, ACS Cent. Sci. 2016, 2, 80.
(14) V. Subramanian, E. E. Wolf, P. V. Kamat, Catalysis with TiO2/gold
nanocomposites. effect of metal particle size on the fermi level equilibration, J. Am.
Chem. Soc. 2004, 126, 4943.
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(15) T. R. Gordon, M. Cargnello, T. Paik, F. Mangolini, R. T. Weber, P. Fornasiero,
C. B. Murray, Nonaqueous synthesis of TiO2 nanocrystals using TiF4 to engineer
morphology, oxygen vacancy concentration, and photocatalytic activity, J. Am. Chem.
Soc. 2012, 134, 6751.
(16) A. Janotti, J. B. Varley, P. Rinke, N. Umezawa, G. Kresse, C. G. Van de Walle,
Hybrid functional studies of the oxygen vacancy in TiO2, Phys. Rev. B 2010, 81,
085212.
(17) C. D. Valentin, G. Pacchioni, A. Selloni, Reduced and n-type doped TiO2: nature
of Ti3+ species, J. Phys. Chem. C 2009, 113, 20543.
(18) Q. Kang, J. Cao, Y. Zhang, L. Liu, H. Xu, J. Ye, Reduced TiO2 nanotube arrays
for photoelectrochemical water splitting, J. Mater. Chem. A 2013, 1, 5766.
(19) M. Setvín, U. Aschauer, P. Scheiber, Y.-F. Li, W. Hou, M. Schmid, A. Selloni, U.
Diebold, Reaction of O2 with subsurface oxygen vacancies on TiO2 anatase (101),
Science 2013, 341, 988.
(20) X. Pan, M.-Q. Yang, X. Fu, N. Zhang, Y.-J. Xu, Defective TiO2 with oxygen
vacancies: synthesis, properties and photocatalytic applications, Nanoscale 2013, 5,
3601.
(21) O. Bikondoa, C. L. Pang, R. Ithnin, C. A. Muryn, H. Onishi, G. Thornton, Direct
visualization of defect-mediated dissociation of water on TiO2(110), Nat. Mater. 2006,
5, 189.
(22) X. Cui, Z. Wang, S. Tan, B. Wang, J. Yang, J. G. Hou, Identifying hydroxyls on
the TiO2 (110)-1×1 surface with scanning tunneling microscopy, J. Phys. Chem. C
2009, 113, 13204.
(23) N. G. Petrik, Z. Zhang, Y. Du, Z. Dohnálek, I. Lyubinetsky, G. A. Kimmel,
Chemical reactivity of reduced TiO2(110): the dominant role of surface defects in
oxygen chemisorption, J. Phys. Chem. C 2009, 113, 12407.
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(24) R. A. Bennett, P. Stone, N. J. Price, M. Bowker, Two (1 × 2) reconstructions of
TiO2(110): surface eearrangement and reactivity studied using elevated temperature
scanning tunneling microscopy, Phys. Rev. Lett. 1999, 82, 3831.
(25) M. Blanco-Rey, J. Abad, C. Rogero, J. Mendez, M. F. Lopez, J. A. Martin-Gago,
P. L. de Andres, Structure of rutile TiO2 (110)-(1 × 2): formation of Ti2O3 quasi-1D
metallic chains, Phys. Rev. Lett. 2006, 96, 055502.
(26) K. T. Park, M. H. Pan, V. Meunier, E.W. Plummer, Surface reconstruction of TiO2
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62
Chapter 4 Activating titania for electrocatalysis by surface vacancy engineering
4.1 Introduction
As an alternative to fossil fuels, hydrogen is regarded as one of the most promising
key energy carriers for a global-scale sustainable energy system.[1] In particular, water-
alkali electrolyzers for overall water splitting exhibit tremendous potential for the
evolution of high-purity hydrogen and oxygen gases, with the additional values of
simple processes, zero CO2 emission and low pollutant. Finding efficient and low-cost
noble-metal-free electrocatalysts towards HER and OER is urgently required for their
large-scale commercial application.[2-6] Thanks to extensive efforts, earth-abundant
transition TMOs have recently been reported to be promising candidates for
electrochemical water splitting in alkaline conditions through defect engineering.[7-11]
These defects are expected to effectively improve the intrinsic electrical conductivity
of defective TMOs, act as catalytic active cites and enhance the catalytic activity.
Among various defects, surface OV is regarded as one of the most important, and
supposed to be the prevalent defect in many TMOs.[12-15] Previous studies have
demonstrated that when oxygen atoms are removed from the TMOs, the geometric,
physical and chemical properties can be profoundly modified. For example, the
coexistence of different oxidation states of metals, such as Ti3+-Ti4+ in TiO2, Co2+-Co3+
in Co2O3 and Ce3+-Ce4+ in CeO2, are essential in the electron translocations and
extremely important for many of their catalytic applications.[9,15,16] Additionally, the
surface OV has been well-recognized as favorable adsorption for adsorbates (such as
H2O, O2, metal nanocluster, CO2, etc) for the defective TMOs in catalytic processes.[17-
22] Obviously, elucidating the role of OV, as well as other defects, in the modified
electrocatalytic properties of defective TMOs is of immense scientific and
technological importance towards an in-depth understanding and optimizing catalysts.
63
However, it remains challenging owing to the lack of detailed knowledges of catalytic
active sites and the electron translocations at the atomic level. Meanwhile, the precise
control of defects is very difficult, because complicated defect compositions along with
a wide diversity of sample conditions, such as morphology, crystallinity, dimension
and diameters are common in defective TMOs.
Here, in this paper, surface OVs were introduced into the highly pure and
stoichiometric TiO2 rutile(110) single crystal, which is a prototypal and widely studied
TMO, through thermal treatment under UHV condition.[23,24] STEM and STM
techniques found that the as-treated single crystal kept a good crystalline structure. In
addition, OVs were found mainly exist in the surface and subsurface region with a
depth of around 100 nm rather than the inner bulk region. The simple sample conditions
allow us to have an in-depth visit on the effect of OV on the physical and chemical
properties of TiO2, as well as its roles in surface electrocatalytic processes. In-situ
synchrotron XPS confirmed the appearance of mid-gap defect state in the as-treated
single crystal, which was aroused by the associating Ti3+ ions with OVs. The catalytic
activities of TiO2 were investigated by electrochemical HER in alkaline conditions,
suggesting a tremendous enhancement for the as-treated single crystal with OVs.
Further STM and DFT calculations suggested Ti3+ ions neighboring at OV could be
helpful for the proton reduction process and promote the electrons translocation. Gibbs
free energy calculations verified that the existence of OV on the surface or subsurface
could effectively low the energy of hydrogen adsorption and benefit the HER.
4.2 Experimental sections
The cross-sectional TEM images were acquired with a STEM (JEOL JEM-
ARM200f) both in HAADF mode and annular bright-field (ABF) mode. In-situ EELS
64
was used to characterize the OV concentration of the cross-sectional sample in TEM
measurements. The cross-sectional sample was prepared by the FIB (Zeiss Auriga FIB-
SEM) technique.
In-situ XPS characterizations were carried out at Beamline 4B9B in the Beijing
Synchrotron Radiation Facility (BSRF), and variable photon energies were referenced
to a fresh Au polycrystalline film. The spot size of incident light in XPS was about 1
mm in diameter. All the data were recorded in UHV at room temperature. The Hall
coefficient and magnetoresistance were measured by the five-probe technique using a
Quantum Design Physical Property Measurement System (PPMS)-14T. The electrolyte
was 1 M KOH solution. The electrochemical HER experiments were performed by a
typical three-electrode method, in which a Pt plate and Hg/HgO (0.923 V versus the
standard hydrogen electrode) were used as the counter and reference electrodes,
respectively. All electrochemical measurements were performed with a Bio Logic
Science Instruments VSP-300 electrochemistry workstation. The linear portions of
Tafel plots were fitted to the Tafel equation: η = blog|J| + a, where η is the overpotential,
a is the exchange current density, and b is the Tafel slope.
All DFT calculations were performed using the Vienna Ab Initio Simulation
Package (VASP). The generalized gradient approximation (GGA) was applied to treat
the exchange correlation energy with the Perdew-Burke-Ernzerhof (PBE) functional.
The projector augmented wave (PAW) method was employed to describe electron−ion
interactions, with the cut-off energy of 400 eV. The structural model of the TiO2(110)
surface was constructed with four Ti-O layers as a 4 × 2 periodic supercell comprising
192 atoms with a vacuum spacing of 20 Å to avoid interaction between adjacent
surfaces. Spin-polarized local density approximation plus on-site Coulomb self-
interaction potential (LDA+U) calculations were performed for the Hubbard
65
correction, and an effective U (Ueff) value of 4.2 eV was applied in all calculations. All
structures in the calculations were relaxed until the convergence tolerance of the force
on each atom was smaller than 0.02 eV. The energy convergence criterion was set to
be 1 × 10-4 eV for self-consistent calculations, and k-point sampling was restricted to
the Gamma point only because of the large size of the supercell.
4.3 Characterization of OV by STEM
Figure 4.1 Crystal structure mode of the TiO2 rutile(110) surface, surf-OV can be
created through removing Obr atoms.
Figure 4.2 SEM image of reduced TiO2 single crystal sample cut by a focused ion beam
(FIB); inset is an enlarged image of the selected area.
66
First of all, pure and stoichiometric TiO2(110) single crystal can be well described
by the structure model shown in Figure 4.1, with alternating rows of bridging oxygen
(Obr) atoms and five-coordinate titanium (Ti5c) atoms lying in a plane along the [001]
direction. The reduced single crystal was obtained through annealing in UHV
conditions with base pressures in the low 1 × 10-10 torr regime at 900 K. The cross-
section of the reduced TiO2 single crystal was characterized by STEM, with the sample
prepared by FIB shown in Figure 4.2. The top surface, near-surface and inner bulk
region all exhibited high crystallinity, shown in Figure 4.3 (a) and (b), indicating that
the formation of OVs did not interrupt its well-crystalline structure of the reduced TiO2.
It should be emphasized that the good crystallinity of the surface area of the reduced
sample could exclude the existence of disorder or amorphous surface layer and foreign
atoms, which are widely studied in many other cases.[25-27] Moreover, the gradually
darkening contrast from inner region to top surface with a thickness of around 80 nm
in the STEM image implied the increasing oxygen deficiency and the displacement of
Ti atoms, because the contrast is proportional to the average atomic number of atomic
columns in HAADF mode.[28]
67
Figure 4.3 Cross-sectional TEM image of reduced TiO2(110) single crystal in ABF,
showing a high crystalline structure both in the surface and bulk regions. (b) Large-
region cross-sectional STEM image of reduced TiO2(110) single crystal in HAADF
mode, in which three areas with different depth from the surface (10 nm, 50 nm, and
100 nm) are marked. (c) and (d) Corresponding in-situ EELS Ti-L edge and low-level
spectra of the three regions marked in (b).
To further reveal the effect of OVs on the reduced TiO2 single crystal, in-situ EELS
were acquired followed the STEM measurement at three different regions (10 nm, 50
nm and 100 nm from top the surface, marked in Figure 4.3 (b). In the Ti-L edge EELS
spectra, the intensity ratio of dz2/dx2-y2 which is sensitive towards the Ti-O bonding
lengths[29] in the top surface region apparently increased compared with those in the
inner regions, agreeing well with the OV induced lattice distortions recognized in
above experiments. As shown in the low-level spectrum in Figure 4.3 (d), the three
68
peaks (6.4 eV, 11.3 eV and 14.6 eV), originating from O 2p orbitals to Ti 3d orbitals,
are marked in the region A. Compared with inner region, the peaks at the top surface
region (10 nm) exhibited obvious broadening, which were caused by the OV induced
lattice distortion and indicated a higher OV concertation near the top surface area.[30]
Meanwhile, the peaks at 25.2 eV (peak B) and 48.5 eV (peak C) in the inner region
shifted towards lower energy to 24.6 eV and 48.2 eV in the top surface region, due to
the emerging of more OVs. Peak B was assigned to transitions between O 2p state and
Ti 4sp states. Peak C was assigned to excitations from the Ti 3p core level to 3d excited
states.[31] Therefore, the existences of OVs are expected to modify the valence state of
their surrounding Ti atoms.
4.4 Characterization the electronic structure by in-situ synchrotron XPS and
UPS
To investigate the effect of OV on the electronic structure of reduced TiO2 single
crystal, in-situ synchrotron XPS spectra and valance band spectra, was carried out. As
shown in Figure 4.4 (a), after annealing the stoichiometric pure single crystal at 900 K
in HUV for 5 h, a broad peak at around 457 eV belongs to Ti3+ ions appeared in the Ti
2p XPS spectra.[32] It means the excess electrons from the removed oxygen atoms were
trapped by Ti atoms and formed Ti3+ ions. In the valence band (VB) spectra shown in
Figure 4.4 (b), a peak, locating about 0.8 eV below the Fermi Level appeared after
annealing, which is known as the Ti 3d defect state. Meanwhile, a broad peak locating
at 12 eV can be seen below the Fermi Level for both samples, which can be assigned
to the OHs.[33] The broad peaks reflected the dissociation of water at the surface, which
is consistent with the STM observation.
69
Figure 4.4 (a) In-situ Ti 2p XPS spectra of the pristine and OV-TiO2 single crystal. (b)
VB spectra of pristine and OV-TiO2 single crystal.
Figure 4.5 DFT calculations of the mid-gap states caused by OVs in TiO2(110). (a)
Calculation of the excess electron distribution of the reduced TiO2(110) surface with
an OV on the top surface. (b) The calculated band structure (left) and DOS (right) of
the TiO2(110) with a surface OV.
In addition, DFT calculation was used to verify the OV-induced changes of
electronic structure in the reduced TiO2. As shown in Figure 4.5, after introducing an
OV on the surface of TiO2(110), the two excess electrons belonged to the removed
oxygen atom will bond with two neighboring Ti atoms and form Ti3+ ions. The other
two cases, including OVs at inner region and surface OHs from water dissociation at
70
surface OV sites were also included in the DFT calculations, which demonstrated that
they exhibited similar results of creating Ti3+ ions (Figure 4.6). As a sequence, defect
states in the gap, can be observed in the calculated band structure and DOS, which are
mainly contributed to the Ti 3d orbitals. Above XPS and DFT results indicate that the
band gap Ti 3d defect states originate from these intrinsic defects and associating Ti3+
ions in the reduced TiO2.
Figure 4.6 Calculation of the excess electron distribution on the TiO2(110) surface with
an OV on the sublayer and with two surface Ad-H. (a) and (b) The calculated band
structure and DOS of TiO2(110) with sublayer OVs, respectively. (c) and (d) The
calculated band structure and DOS of TiO2(110) with surface Ad-Hs, respectively.
71
4.5 Electrochemical HER activity of OV-TiO2
Figure 4.7 (a) LSV data on different electrocatalysts for the HER at the rate of 10 mV∙s-
1, inset shows the pristine TiO2(110) single crystal (P), OV-low TiO2(110) single
crystal (L), OV-high TiO2(110) single crystal (H), and the Nb-doped TiO2(110) single
crystal (Nb). The OV-high TiO2 was tested for 1000 cycles with the LVS curve shown
as the red solid and green dashed lines for before and after cycling, respectively. (b)
Cycling stability over 18 h of the OV-high sample at a potential of –0.7 V. (c) Tafel
plot of the OV-high TiO2. (d) Cdl determined from the linear fitting of the capacitive
current vs. scan rate, measured from the cyclic voltammetry (CV) curves of OV-high
TiO2, OV-low TiO2 and Nb-doped TiO2.
To evaluate the electrochemical HER activity of the reduced TiO2 single crystal, two
reduced TiO2 single crystal annealed in HUV at 900 K for 5 h (OV-low) and 50 h (OV-
high), and a reference sample (Nb-doped (0.43 at%)) single crystal were selected. The
electrochemical behavior for the HER of these samples were then tested at 1 M KOH
aqueous solution with a scan rate of 10 mV/s. As shown in Figure 4.7, OV-high TiO2
72
exhibited considerable HER activity with a current density of 22.8 mA.cm-2 at -0.80 V,
while Nb-doped TiO2 exhibited very weak activity and OV-low TiO2 was almost not
active under the same experimental conditions (IR compensation based on resistance
test was applied for all linear sweep voltammetric (LSV) curves, Figure 4.8). Long-
term stability of the OV-high TiO2 was demonstrated by the cycling test of 1000 times
(red dash line in Figure 4.7 (a)), and a continuous test at a potential of –0.7 V for more
than 18 h (Figure 4.7 (b)).
The linear portions of Tafel plots were fitted to the Tafel equation: η = blog|J| + a,
where η is the overpotential, a is the exchange current density, and b is the Tafel slope.
The Tafel slope of OV-high TiO2 near the substantial cathodic current region from the
HER was plotted to be 187.5 mV.dec-1, as shown in Figure 4.7 (c), which represents
the hydrogen generation rate with the applied overpotential. Figure 4.7 (d) shows the
double layer capacitance (Cdl) of the three samples, which were valued by fitting the
slope of the capacitive current vs. scan rate measured from the cyclic voltammetry (CV)
curves (Figure 4.9). The Cdl value has a proportional relationship with the
electrochemically active surface area (EASA) of catalysts. Therefore, the Cdl value of
OV-high TiO2 at 72 mF.cm-2 demonstrated an apparent increase in the EASA compared
with Nb-doped TiO2 with a Cdl value of 8.8 mF.cm-2, which match well with their HER
performances. It should also be noticed that despite the Cdl value of OV-low (16.0
mF.cm-2) is higher than that of Nb-doped TiO2, its HER performance was severely
suppressed by its low conductivity, which will be presented in Figure 4.11. Therefore,
it is suggested that the elevated electrochemical HER of TiO2 not only originate from
the increased intrinsic conductivity and the charge carrier density, but also the modified
intrinsic electrocatalytic properties that needed to be further clarified.
73
Figure 4.8 Electrochemical impedance spectra of different electrodes at −0.3 V versus
eversible hydrogen electrode (RHE) (inset is the full range measurement).
Figure 4.9 CV curves of different samples under scanning rate ranging from 20 mV/s
to 180 mV/s.
4.6 Determination of the origin of the electrochemical HER activity of OV-TiO2
4.6.1 Electrical conductivity
Electrical conductivity and the number of electrocatalytic sites are considered to be
of great significance for the catalytic performance of electrocatalysts. First of all, five-
probe Hall effect measurements were carried out to reveal the conductivities of all the
74
TiO2 single crystals, as shown in the five probes configuration used in the PPMS
measurement in Figure 4.10.
Figure 4.10 The holder used in PPMS measurement and the five probes configuration
used in the PPMS measurement.
The electron density of n-type semiconductors has an inversely proportional
relationship with the Hall coefficient (RH) as n = ‒1/(e∙RH). As shown in Figure 4.11,
OV-high TiO2 had an electrical resistivity of 80.0 × 10-3 Ω∙m that is lower than that of
the OV-low TiO2 (94.0 × 10-3 Ω∙m). The charge carrier density in OV-high TiO2 (1.9
× 1017 cm-3) is about 40 times higher than that in the OV-low TiO2 (4.8 × 1015 cm-3).
The conductivities of both reduced TiO2 samples are believed to arise from the mid-
gap state induced by OVs. The reference Nb-doped TiO2 single crystal exhibits the
highest electrical conductivity of 40.6 × 10-3 Ω∙m and the highest charge carrier density
75
of 3.2 × 1017 cm-3 among all the samples. The weak electrocatalytic activity of the
reference Nb-doped TiO2 single crystal, however, demonstrates that good electrical
conductivity is inadequate to enable TiO2 activity towards the HER in alkaline media.
Figure 4. 11 Hall resistivity measurements of different samples. The Hall coefficient
RH was determined by fitting the slope of the curve of the Hall resistivity vs. magnetic
field.
4.6.2 Electrochemical activate sites
In order to reveal the catalytic process associated with active sites at the atomic level,
STM investigations were carried out on the OV-high TiO2(110) single crystal. As
shown in Figure 4.12, all the adsorbed H2O molecules reside on OV sites. as confirmed
by the adsorption dynamics observed in Figure 4.12 (a) and (b). OHs also appeared on
the OV sites after dissociation of H2O molecules, as observed in the dynamic processes
in Figure 4.13. This phenomenon suggests that surface OVs are highly active towards
adsorbing and dissociating residual H2O molecules in UHV.
76
Figure 4.12 In-situ STM studies on electrocatalytic dynamics occurring on TiO2(110)
surface associated with oxygen vacancies. (a) STM image of the partially hydroxylated
TiO2 surface, with the OVs indicated by light blue arrows (15 nm × 15 nm, 1.2 V, 20
pA). (b) STM image of the same region of (a), with the OVs mostly filled by H2O (15
nm × 15 nm, 1.2 V, 20 pA).
77
Figure 4.13 (a) STM image of an individual Ad-H2O at an OV site, with another OV
included as a reference point (3 nm × 6 nm, 1.2 V, 10 pA). (b) STM image of the two
OHs from the dissociation of Ad-H2O at the OV site (1.2 V, 10 pA). (c) STM image of
the same area in a and b, in which an OH was removed through a 2.5 V pulse by the
STM tip (1.2 V, 10 pA).
78
Figure 4.14 (a) STM image of the reduced TiO2 surface with two individual OVs in the
empty-state (1.2 V, 20 pA). (b) STM image of the same region of (a), but in the filled-
state, with Ti3+ ions indicated by red arrows (-2.3 V, 10 pA).
In addition, each OV induces two surrounding Ti3+ ions, which was revealed in STM
images taken with negative sample bias (filled-state), as shown in Figure 4.14. Recent
investigations pointed out that the Ti3+ ions in reduced TiO2 exhibit polaron
behavior.[34-36] This enables Ti3+ ions to rapidly hop across the nearby lattice in rutile,
and thus, leads to the increase of the conductivity of the reduced rutile TiO2. Since the
electrical conductivity of reduced TiO2 is proportional to the concentration of Ti3+ ions,
this supports the proposition that the OV-high TiO2(110) single crystal possesses
higher conductivity than the OV-low sample. In order to simulate electrocatalytic HER
processes, we applied a sample bias to allow the STM tip to act as an electron donator
to the OV-high TiO2(110) single crystal surface. Both Ad-H2O and OHs disappeared
from the sample surface when the bias was higher than the threshold of 2.0 V, as shown
in Figure 4.15 (a) and (b). Meanwhile, the OVs at corresponding sites were also healed.
This indicates that hydrogen desorption of Ad-H2O and OHs occurred on the OV-high
TiO2(110) surface at the cost of consumption of OVs. In this case, the HER activity of
79
reduced TiO2(110) single crystal is expected to be depressed gradually with a
decreasing concentration of OVs. Nevertheless, this contradicts our cycling stability
results, in which the OV-high TiO2(110) surface exhibited excellent stability and
durability, suggesting that OVs in the surface are not adequate to drive the
electrocatalytic HER constantly and steadily, although they are active towards
adsorbing H2O molecules and OHs.
Figure 4.15 STM image of the partially hydroxylated TiO2 surface, with OVs, OHs,
and Ad-H2O appearing on the surface (1.2 V, 20 pA). (f) STM image of the same region
of e after the OHs and Ad-H2O were removed by the last scan with a tip bias of 2.5 V.
(1.2 V, 20 pA).
80
4.6.3 DFT calculation of the Gibbs free energy
The origins of the electrocatalytic HER activity of the reduced TiO2(110) surface in
alkaline media was further revealed by DFT calculations, in terms of thermodynamics
and kinetics. The surface structures were modeled according to STEM and STM results
obtained on TiO2(110) surface with OVs.
The calculation methods for the hydrogen evolution reaction (HER) in alkaline
solutions are summarized as follows:
By considering the standard hydrogen electrode as the reference potential and
hydrolysis of water in the solution, the free energy of reactions (4.1) and (4.2) is set to
zero because they are in equilibrium. Therefore, the free energy of (H+ + e-)
corresponds to that of ½ H2 (1 bar, 298 K) and the free energy of (OH- ‒ e-) is calculated
according to the free energy of H2O and (H+ + e-),
(H+ + e-) → ½ H2 (4.1)
H2O → (H+ + e-) + (OH- ‒ e-) (4.2)
We used gas-phase H2O at 0.035 bar as the reference state, because at this pressure,
gas-phase H2O is in equilibrium with liquid water at 300 K. The calculated free
energies of the H2O, (H+ + e-), and (OH- ‒ e-) are listed in Table 4.1.
The Gibbs free energy of the intermediates were calculated as[37]
ΔG = ΔE + ΔZPE ‒ TΔS (4.3)
where ΔE is the binding energy of intermediates which is defined as the reaction
energies of the reactions
H2O + * → OH* + H* (4.4)
H2O + * → OH* + ½H2 (4.5)
½H2 + * → H* (4.6)
81
ΔZPE and ΔS can be obtained from the vibrational frequency ʋi, which are the
changes in zero point energies (ZPE) and entropies due to the reaction, respectively.
All the parameters have been taken from DFT calculations.
Zero point energies are calculated as follows,
ZPE = ∑i ½hʋi (4.7)
i = 3n‒5 (for linear molecule)
i = 3n‒6 (for non-linear molecule)
where i is the degree of freedom for the molecule, and n is the number of atoms in
the molecule.
Entropies are calculated from the sum of the translational entropy St, the rotational
entropy Sr, and the vibrational entropy Sv as follows:
S = St + Sr + Sv (4.8)
(4.9)
(for linear molecule) (4.10-1)
(for non-linear molecule) (4.10-2)
(4.11)
where kB, NA, and h are the Boltzmann constant, Avogadro constant, and Planck
constant, respectively. N, m, V, T, I, and σ are the number of particles, and the mass,
volume, temperature, moment of inertia, and symmetry number of the molecules,
respectively. Otherwise, the entropies of intermediates are calculated by the vibrational
entropy , because no translational or rotational behaviors can be found for an
adsorbed molecule. The calculated ZPE and S of free H2O, H2, and (OH- - e-) are listed
in Table 4.1.
82
At a pH different from 0, we can correct the free energy of H+ ions by the
concentration dependence of the entropy: ΔGpH = ‒kT∙ln[H+] = kT∙ln10∙pH. In our
work, pH is 14 and ΔGpH = 0.83 eV.
The effect of a bias ΔGU was imposed on each step by including an electron in the
electrode as an ‒eU term, where U is the electrode potential relative to the standard
hydrogen electrode. Therefore, the reaction free energy of processes was calculated as:
ΔG(U, pH) = ΔG + ΔGpH + ΔGU (4.12)
As is shown in the reactions of the HER in alkaline solution,
Volmer H2O + e- → H* + OH- (4.13)
Heyrovsky H2O + e- +H* → H2 + OH- (4.14)
Tafel 2H* → H2 (4.15)
We include the effects of pH and U on steps involving (OH- - e-). The free energies
of possible steps (such as water splitting, hydroxide adsorption, hydroxide desorption,
and hydrogen production) were calculated and compared to find out the most optimal
path in our calculation. As a result, the reactions in HER and corresponding free
energies under applied potential U and pH can be written as:
Volmer 1 - water splitting
H2O + M → H* + OH*
ΔG1 = ΔG (H* + OH*)
Volmer 2 - OH* desorption
OH* → (OH- ‒ e-) + M
ΔG2 = ‒ΔG(OH*) + eU + 0.83
Heyrovsky
H2O + H* → H2↑ + (OH- ‒ e-) + M
ΔG3 = ‒ΔG(H*) + eU + 0.83
83
Tafel
2H* + M → H2↑
ΔG4 = ‒ΔG(2H*)
Here, M represents the catalyst, which can be a TiO2(110) surface with surfOV and
subOV, only subOV, or subOV with Obr-H surface. The calculation results of main
reaction pathway are listed in Table 4.1.
Table 4.1. Binding energies, entropies, and zero point energies contribution to Gibbs
free energy of main reaction pathway. subOV + Obr-H* is regard as M.
subOV + Obr-H* ΔE ZPE ΔZPE TS TΔS
ΔZPE
- TΔS
ΔG
ΔG
(pH)
H2O 0.566 0 0.534 0 0 0 0
½ H2 0.133 0.201
OH- ‒ e- 0.433 0.333
H2O* -0.932 0.631 0.065 0.020 -0.514 0.579 -0.353 -0.353
H* + OH* -0.735 0.538 -0.028 0.009 -0.525 0.497 -0.238 -0.238
H*+ (OH- ‒ e-) -0.600 0.671 0.105 0.333 -0.201 0.306 -0.294 0.536
½H2 + (OH- ‒ e-) 0.566 0 0.534 0 0 0 0.83
84
Figure 4.16 DFT calculations reveal roles of oxygen vacancies in electrocatalysis on
TiO2(110) surface. (a) Free energy pathways of the relevant reaction intermediates in
alkaline media on the reduced TiO2 with both surfOV and subOV, with the OVs
marked as solid black circles. (b) Compare of HER free energy of pristine TiO2 and
reduced TiO2 with subOV and Obr-H* (an onset potential of URHE = ‒1.15 V was applied
for both catalysts).
Therefore, based on our calculated, the Free energy pathways of the relevant reaction
intermediates in alkaline media on the OV-TiO2 with both surfOV and subOV, is
shown in Figure 4.16. It was found that water molecules prefer to be dissociated at
surface OV (surfOV) site in initial stage (step 1). The produced hydrogen atoms are
favorable towards combing in pairs and forming hydrogen molecules (step 2 to 3). All
the reaction steps are exothermic processes. This supports our hypothesis that the
surface OVs are easily to be healed during the electrochemical reaction. Hence, it
indicates that surfOVs were not sustainable which agrees with STM. HER pathway on
the reduced TiO2 with sublayer OV (subOV), as well as the pristine TiO2 were further
calculated, as shown in Fig. 4.16 (b). Hydrogen desorption reaction (Heyrovsky
85
reaction) step between intermediate 2 and intermediate 3 is deter determined to be the
rate-limiting step with the largest free energy difference for both pristine TiO2 and
reduced TiO2. The pristine TiO2 needs a theoretical potential of URHE = ‒1.33 eV to
overcome this energy barrier. In contrast, the potential barrier can be effectively
lowered on the reduced TiO2 to enable all HER reaction steps to be exothermic and
energetically favorable, due to the presence of subOV. This has been confirmed
experimentally in our electrocatalytic HER process, in which onset potential for HER
in reduced TiO2 was significantly decreased. Therefore, the theoretical calculations
suggest that the electrocatalytic HER activity of reduced TiO2(110) single crystal
mainly originates from subOVs, which can effectively promote the hydrogen
desorption capability and lower the overpotential of HER in alkaline media.
Figure 4.17 Optimized pathway of the HER in alkaline media on the reduced TiO2
with surfOv and subOV based on the free energy calculation.
Figure 4.17 gives the detailed optimized overall pathway of the HER in alkaline
media on the OV-TiO2 with surfOv and subOV based on the free energy calculation,
with all the intermediates listed below:
Intermediates
1: TiO2-surfOV-subOV + H2O TiO2-subOV + surfOV-H2O*
2: TiO2-subOV + surfOV-H2O*
86
3: TiO2-subOV + Obr-H* + Ti-H*
4: TiO2-subOV + H2
5: TiO2-subOV + Ti5C-H2O*
6: TiO2-subOV + Ti5C-OH* + Obr-H*
7 and 11: TiO2-subOV + Obr-H* + OH-
8 and 12: TiO2-subOV + 2Obr-H* + Ti5C-OH*
9 and 13: TiO2-subOV + 2Obr-H* + OH-
10 and 14: TiO2-subOV + Obr-H* + OH- + H2
15: TiO2-subOV + H2
All reactions are written as:
a. TiO2-surfOV-subOV:
Volmer (1-2-3)
TiO2-surfOV-subOV + H2O → TiO2-subOV + surfOV-H2O*
TiO2-subOV + surfOV-H2O* → TiO2-subOV + Obr-H
* + Ti5C-H*
Tafel (3-4)
TiO2-subOV + Obr-H* + Ti5C-H*→ TiO2-subOV + H2↑
b. TiO2-subOV:
Volmer (4-5-6)
TiO2-subOV+ H2O → TiO2-subOV + Ti5C-H2O*
TiO2-subOV + Ti5C-H2O* → TiO2-subOV + Ti5C-OH* + Obr-H
*
Heyrovsky (6-11)
TiO2-subOV + Ti5C-OH* + Obr-H* → TiO2-subOV + Obr-H
* + (OH- ‒ e-)
c. TiO2-subOV + Obr-H*:
Volmer (11-12-13)
TiO2-subOV + Obr-H* +H2O → TiO2-subOV + 2Obr-H
* + Ti5C-OH*
87
TiO2-subOV + 2Obr-H* + Ti5C-OH* → TiO2-subOV + 2Obr-H
* + (OH- ‒ e-)
Heyrovsky (13-14)
TiO2-subOV + 2Obr-H* + H2O → TiO2-subOV + Obr-H
* + H2↑ + (OH- ‒ e-)
Tafel (13-15)
TiO2-subOV + 2Obr-H* → TiO2-subOV + H2↑
4.7 Summary
Our study shows that inactive pure TiO2 rutile(110) single crystal can be activated
towards the HER in alkaline media through creating OVs and accompanying Ti3+ ions
by annealing in UHV. OVs and Ti3+ ions in the surface region dominate the electrical
conductivity of reduced TiO2 and the amount of electrocatalytic active sites.
Combining the well-characterized atomic surface structure and theoretical
calculations, we conclude that subOVs can promote the electron transfer and hydrogen
desorption in the electrocatalytic HER on reduced TiO2 in alkaline media. Considering
that all the electrochemical characterizations were performed on a single crystal sample
with atomically flat surface and low specific surface area, the overpotential is expected
to be significantly decreased in reduced TiO2 nanoparticles with a higher density of
active sites. Our work helps to elucidate the fundamental mechanism of the
electrocatalytic activity of reduced oxides towards the HER, which is of immense
fundamental and practical importance towards an in-depth understanding and rational
optimization of TMOs as electrocatalysts in alkaline media.
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Chapter 5 Tuning electronic structure of BiOBr 2D nanosheets by strain for
photocatalysis
5.1 Introduction
The ability to continuously control the electronic structures in photocatalysts is
highly desirable for a wide range of energy and environmental applications, including
H2 production by water splitting, carbon fixation, and the elimination of pollution.[1,2]
For example, the absorption spectrum and the quantum conversion efficiency of a
photocatalyst, which determine its performance, can be modulated by tuning the band
gap (Eg), the positions of the VB and the conduction band (CB), and the band dispersion
of photocatalysts.[3] In a similar way to chemical composition, strain is a continuous
variable that is capable of altering electronic structure. Although strain engineering is
a straightforward method, its potential in photocatalysis remains largely under-
exploited.
Bismuth oxyhalides BiOX (X = Cl, Br, I) are p-block photocatalysts that have
attracted considerable attention due to their unique 2D layered structure and excellent
photocatalytic properties under visible light.[4-9] In BiOX, [Bi2O2]2+ slabs are
interleaved with double halogen atoms slabs by strong electrovalent bonds along the
[001] direction, while two closely adjacent slabs of halogen atoms are connected by
van de Waals interactions.[10,11] The dispersive VB and CB induced by sp hybridization
give rise to high mobility of the photo-induced charge carriers. In addition, the internal
electric field resulting from the asymmetric charge distribution between [Bi2O2]2+ and
the halogen layers facilitates the effective separation of these photo-induced charge
carriers, and hence, enables the photocatalytic activity of BiOX.[12-14] Due to their 2D
94
layered structure, the electronic structure of BiOX compounds is highly sensitive to
even a subtle inner strain variation.[15] It is, however, a practical challenge to modulate
the photocatalytic properties through fine-tuning the inner strain in nanoscale BiOX.
In this work, we illustrate experimentally and theoretically that the photocatalytic
performance of BiOBr nanosheets can be tuned by the inner strain effect. The
characterizations of photocatalytic degradation and geometric phase analysis (GPA) of
the TEM images indicate that the distribution and intensity of the inner strain dominate
the photocatalytic activity of BiOBr nanosheets. DFT calculations demonstrate that the
strain-modulated photocatalytic properties in BiOBr originate from variation of the
intrinsic electronic structure of this photocatalyst.
5.2 Experimental section
In the synthesis procedure, 5 mmol Bi(NO3)3∙5H2O and 5 mmol cetyl
trimethylammonium bromide (CTAB) were added into 100 mL distilled water at room
temperature with stirring for 20 min, and then, 1 M NaOH solution was added to the
solution to adjust the pH value to 7, 5, 3, and 2 (with the resulting samples denoted as
BiOBr-1, BiOBr-2, BiOBr-3, and BiOBr-4, respectively). The mixed solution was
stirred for 1 h, and then poured into a 100 mL Teflon-lined stainless autoclave up to
80% of the total volume. The autoclave was heated at 170 °C for 17 h, and then cooled
to room temperature in air. The resulting precipitates were collected, washed with
ethanol and deionized water several times, and dried at 80 °C for 10 h.
The morphologies and microstructures of the as-prepared samples were
characterized by SEM (Hitachi CS 3400) and TEM (JEOL JEM-3010, operated at 300
kV). In the photo-degradation experiments, BiOBr powder (100 mg) was added to an
aqueous solution of Rh B (0.02 mmol/L, 100 mL) in a 150 mL quartz reactor.
95
Degradation experiments on Rh. B using N-doped P25 TiO2 nanopowders were also
carried out under the same conditions as a reference. For the photo-degradation of MO
(10 mg/L, 50 mL), 50 mg catalyst was added to an aqueous solution of MO (10 mg/L,
50 mL) in a 150 mL quartz reactor. All photocatalytic experiments were carried out
under irradiation by a 300 W Xe lamp with a filter glass (λ ≥ 420 nm) to remove
ultraviolet (UV) light after dark adsorption experiments. The absorption spectra of Rh.
B and MO were collected on a Hitachi U3010 UV-Visible (UV-Vis)
spectrophotometer. In the photocurrent-time response system, a 300 W Xe lamp with
a monochromator and a cut-off filter (λ ≥ 420 nm) was used as the light source. UV-
Vis diffuse reflectance spectra were collected on a Cintra-10e spectrometer. The
distributions of in-plane strain across the nanosheets were identified by GPA based on
the HRTEM images. In this paper, STEM-CELL was used to simulate the distribution
of the strain, based on the HRTEM images, by calculating and analysing the fast
Fourier transform (FFT) and inverse FFT of the entire images. In the analysis process,
the displacement of lattice parameters (u) are determined by calculating and analysing
the Fourier transform of the selected image. The strain is then obtained from the
derivative of the displacement in the picked direction, and visualized by inverse Fourier
transform.
The first-principles calculations were performed using the VASP code. The GGA
was applied to treat the exchange correlation energy with the PBE functional. The PAW
method was employed to describe the electron-ion interactions. A k-point sampling of
9×9×6 was generated with original Gamma meshes. The cut-off energy for the plane
wave basis was 550 eV. The biaxial strain simulations were realized by fixing the x
and y axes and optimizing the z axis. Equilibrium geometries were obtained by the
minimum energy principle.
96
5.3 Morphology and structure characterization
Figure 5.1 SEM images of (a) BiOBr-square and (d) BiOBr-circle (scale bar is 1 m).
Cross-sectional TEM images of (b) BiOBr-square and (e) BiOBr-circle (scale bar is
200 nm); insets are the HRTEM images of the cross-sections, where both samples show
layer structure along the c-axis (scale bar is 10 nm). HRTEM images of (c) BiOBr-
square and (f) BiOBr-circle, with the corresponding SAED patterns as the insets,
indicating that both samples have the (001) face exposed (scale bar is 10 nm-1 and 5
nm-1 for the insets).
Figure 5.1 shows two typical morphologies of BiOBr nanosheets, square-shaped
(also assigned as BiOBr-1) and circle-shaped BiOBr (BiOBr-4), which were fabricated
by hydrothermal reaction with different pH values. We found that the BiOBr
nanosheets underwent a morphology transition from square-like to circle-like when the
pH value was decreased. The BiOBr nanosheets are several micrometres in size. The
thicknesses of the BiOBr-square and BiOBr-circle nanosheets are 31 nm and 22 nm,
respectively, as revealed by TEM in Figure 5.1 (a) and (b). The insets demonstrate the
layered structure of the BiOBr nanosheets. The interlayer distance is approximate 0.8
nm for both samples. As shown in Figure 5.1 (c)-(f), the d-spacing of 0.27 nm indicates
97
that the (110) face is along the in-plane BiOBr nanosheets. The corresponding insets
show the selected area electron diffraction (SAED) patterns, in which the marked spots
can be indexed as (200) face, (110) face and (1-10) face. Hence, the exposed surfaces
can be identified as {001} facets for both the BiOBr-square and the BiOBr-circle
nanosheets.
5.4 Characterization of strain in BiOBr 2D nanosheets
By carefully examining the X-ray diffraction (XRD) patterns, obvious diffraction
peak shifts could be identified. Figure 5.2 shows the XRD patterns of the as-prepared
BiOBr nanosheets. All the samples (labelled as BiOBr-1 to BiOBr-4) prepared by the
hydrothermal method with different pH values are well-crystallized single-phase
nanopowders. The diffraction patterns can be indexed to the tetragonal structure
(P4/nmm(129)) according to the standard data (PDF card #09-0393). By precisely
controlling the concentration of NaOH, the diffraction peaks can be shifted gradually
in accordance with the pH value of the precursors, as shown in Figure 4.2 (a). In
particular, the (110) peak and the (004) peak demonstrate a clear upshifting with
increasing pH value, as shown in Figure 5.2 (b)-(d). BiOBr possesses a typical 2D
layered crystal structure, in which [Bi2O2]2+ slabs are interleaved with double bromine
atom slabs by strong electrovalent bonds along the [001] direction, as shown in the
insets of Figure 5.2 (c) and (d). The diffraction peak shift indicates the presence of in-
plane strain in BiOBr. The change in the corresponding interplanar crystal spacing,
identified by the refined cell parameters obtained using MDI Jade 5.0 (Materials data,
Inc., Livermore, CA, 1999), was used to estimate the degree of in-plane strain in BiOBr
by comparing it with the theoretical d-value. As shown in Figure 5.2 (c) and (d), in
contrast to the BiOBr-4 sample, which exhibits tiny strain, BiOBr-1 shows a much
98
higher degree of compressive strain for both the {110} and the {001} facets. More
importantly, our results indicate that by carefully controlling the pH value during
sample preparation, the crystal strains were gradually changed for the BiOBr
nanosheets by a facile chemical method.
Figure 5.2 (a) XRD patterns of BiOBr samples fabricated with different pH values.
(b) Local areas of (110) and (004) peaks of BiOBr, where obvious shifts of the peaks
are observed, which are indicated for different BiOBr samples. (c) and (d) Plots of the
changes in the peak positions and the relative tensile strain compared with the
theoretical crystal structure of the (110) and (004) faces, respectively; insets are
schematic illustrations of the BiOBr crystal structure.
99
Figure 5.3 TEM images of (a) BiOBr-square and (c) BiOBr-circle (scale bars are 1000
nm). Strain simulation of (b) BiOBr-square and (d) BiOBr-circle based on HRTEM
(scale bars are 10 nm). The internal strain distributions are in the xy-direction (Exy), the
x-direction (Exx) and the y-direction (Eyy), with the scale for the whole image area
In order to reveal the details of the morphology dependence of the inner strain, we
carried out TEM characterization and GPA simulation based on the HRTEM images,
as shown in Figure 5.3.[16-18] The in-plane wrinkles observed in the TEM images of the
BiOBr-square nanosheets reflect the existence of a large inner strain, while the BiOBr-
circle sample exhibits a relatively strain-free character, as shown in Figure 5.3 (a) and
(c). Figure 5.3 (b) and (d) show the inner strain distribution maps of the BiOBr
nanosheets in the xy-direction (Exy), the x-direction (Exx), and the y-direction (Eyy),
respectively, as obtained by strain simulation. The inhomogeneous compressive strain
100
distribution in the BiOBr-square nanosheets is reflected by the severe local lattice
distortions across the whole surface in Figure 5.3 (b). In contrast, the BiOBr-circle
nanosheets exhibit a quite uniform strain distribution, and this sample shows much less
lattice distortion in the strain maps in Figure 5.3 (d). It should be noted that the strain
difference across the BiOBr-square nanosheets is higher than that across the BiOBr-
circle nanosheets. For example, the strain in the BiOBr-square sample varies from 0.73
to 1.29 (Exy), from 0.78 to 1.21 (Exx), and from 0.80 to 1.16 (Eyy), while in BiOBr-
circle, it varies from 0.88 to 1.17 (Exy), from 0.87 to 1.10 (Exx), and from 0.93 to 1.08
(Eyy). These results confirm that the inner strain of the BiOBr nanosheets can be varied
by the morphology, which can be precisely tuned by the fabrication conditions.
5.5 Characterization of photocatalytic activity of BiOBr 2D nanosheets
Firstly, UV-Vis diffuse reflectance spectra of the BiOBr nanosheets with different
strains were carried out to determine their band gap, which has a significant impact on
their light absorption ability. The band gap of semiconductor can be determined by the
Tauc formula,
(5.1)
(5.2)
Here, α, ν, Eg, A, n, and λ are the absorption coefficient, the incident light frequency,
the band gap, a constant, an integer, and the absorption edge, respectively. Firstly, the
approximate Eg can be calculated by Equation 5.2, and then ln(αhν) vs. ln(hν-Eg) is
plotted; thus, by refining the slope of the straightest line near the band edge we can
obtain the value of n. Secondly, (αhν)2/n vs. hν is plotted, and then the band gap Eg is
evaluated by drawing an extension line to the hν axis intercept. Finally, the accurate
101
value of Eg of BiOBr-square is calculated as 2.68 eV, and 2.82 eV is calculated for
BiOBr-circle. This indicates that the inner strain has a significant effect on the band
gap of BiOBr.
As shown in Figure 5.4, the BiOBr samples exhibit excellent photocatalytic
performance in Rh B degradation. Their photocatalytic activities are higher than for N-
doped P25 TiO2 nanopowders. The low-strain BiOBr-circle sample shows the highest
activity among all the samples, and it demonstrates almost 100 % degradation of Rh.
B within 30 min. As demonstrated in Figure 5.4 (c), the apparent first-order rate
constant (k) for BiOBr-circle is almost twice those for BiOBr-square and the N-doped
P25 TiO2 nanopowders. Similar results were also observed for degradation of other
dyes, such as MO, as shown in Figure 5.4 (d). As shown in Figure 5.4 (e), both the
BiOBr-square and the BiOBr-circle samples can degrade phenol under visible light.
Again, the BiOBr-circle nanosheets show better visible-light photocatalytic
degradation performance on phenol than the BiOBr-square nanosheets. The
photocurrent response of the BiOBr samples were measured for several On-Off cycles
under visible light irradiation. As shown in Figure 5.4 (f), the photocurrent of the
BiOBr samples exhibits a quick response to light irradiation, with the photocurrent
sharply decreasing to zero as soon as the light is turned off, while the photocurrent
quickly reaches stable values when the light is turned on. The stable photocurrents
measured on the BiOBr-square and BiOBr-circle samples under visible light are 0.7
µA and 1.5 µA, respectively. The higher photocurrent of BiOBr-circle than BiOBr-
square under visible light suggests that more efficient photoexcited charge carrier
separation and less recombination of electron-hole pairs were possibly achieved by
decreasing the inner strain in the BiOBr nanosheets.
102
Figure 5.4 (a) UV-Vis diffuse reflectance spectra of BiOBr-square and BiOBr-circle;
inset shows the derivation of the band-gap values for BiOBr. (b) Degradation
experiments on Rh. B by P25, BiOBr-square, and BiOBr-circle under visible light
irradiation, and (c) the corresponding apparent rate constants. (d) Degradation
experiments on MO by BiOBr-square and BiOBr-circle under visible light irradiation,
and (e) Degradation experiments of phenol by BiOBr-square and BiOBr-circle
nanosheets under visible-light irradiation. (f) Current density transient with light on/off
for BiOBr powders under light irradiation.
5.6 DFT of the strain effect on the electronic structure of BiOBr
We carried out DFT calculations in order to the reveal the strain effect (both
compressive and tensile strains) on the electronic structure of BiOBr, which dominates
the photocatalytic properties.[19-21] The strain-free and compressive-strained BiOBr
show the typical electronic features of an indirect-band-gap semiconductor. It changes
to a direct band-gap semiconductor, however, if tensile strain is present in BiOBr. With
a 9.1% tensile strain in the BiOBr lattice, the valence band maximum (VBM) of BiOBr
moves from the R point to the Z point, while the CBM is still located at the same high
symmetry point (Z point) in k-space. It is found that the Eg of BiOBr can be modulated
by the strain effect, as shown in Figure 5.5 (a). For example, the band gap varies
103
between 1.94 eV, 2.12 eV, and 1.27 eV in BiOBr which has 8.8% compressive strain,
is free of strain, and has 9.1% tensile strain, respectively. Both compressive and tensile
strains lead to a narrowed band gap, which is also observed in the other indirect-band-
gap semiconductors.[22] Figure 5.5 (b) shows the calculated DOS. It is found that the
bottom of the CB of strain-free BiOBr is mainly contributed by Bi 6p, and the top of
the VB is dominated by Br 3p, O 2p, and Bi 6s orbitals. While under tensile strain, the
contribution of Bi 6s to the VBM is suppressed.
Figure 5.5 DFT calculations of the band structure of BiOBr with biaxial strain. (a) The
left panel models compressive strain, the central panel is for a strain-free sample, and
the right is for tensile strain. (b) Calculations of the DOS of BiOBr with different kinds
of strain.
Several key factors can affect the photocatalytic activity of BiOBr nanosheets in
photocatalytic reactions, which include photon absorption, separation of photoexcited
carriers, and surface area. Our UV-Vis diffuse reflectance spectra results demonstrate
that the band gap of BiOBr-square (Eg = 2.68 eV) is smaller than that of BiOBr-circle
(Eg = 2.82 eV). BiOBr-square is, therefore, expected to have a broader range of light
absorption, in contrast to BiOBr-circle. BET measurements reveal that the BiOBr-
square sample exhibits a larger surface area of 7.03 m3/g compared to BiOBr-circle
(4.49 m3/g) (Table 5.1). It is interesting, however, that the photocatalytic measurements
indicate that the photocatalytic activity of BiOBr-sqaure is much lower than that of
BiOBr-circle. This is because the photoexcited charge separation in BiOBr determines
104
the overall photocatalytic activity in the photocatalytic process, which was also
reported in previous works.[23,24] As shown in Figure 5.4 (f), the photocurrent
measurements suggest that the charge separation is indeed more efficient in BiOBr-
circle nanosheets than in BiOBr-square nanosheets. Based on the DFT calculation
results, we believe that the electronic structure of the BiOBr-circle sample may
facilitate the separation of photoexcited charge carriers. It was found that the
interactions between Br atoms and [Bi2O2]2+ slabs mediated by weak van de Waals
coupling are tunable by the strain effect. In other words, the band symmetry of BiOBr
can be modulated by strain through tuning the electronic interaction between the Br
atoms and the [Bi2O2]2+ slabs. The dramatic changes in the band symmetry, e.g. from
direct to indirect band gap or the change of energy dispersion due to the strain, would
affect the separation of photoexcited charge carriers. The electron-hole recombination
occurring in indirect semiconductors typically requires the emission of multiple
phonons to accommodate the energy and momentum differences between the CB and
the VB. An appropriate rearrangement of the electronic symmetry, for instance, in the
BiOBr-circle sample, may tune the momentum mismatch and improve electron-hole
separation. Moreover, the VB and CB in strain-free BiOBr are more dispersive, which
is expected to lead to a high mobility of photoexcited charge carriers. In the case of the
BiOBr-square nanosheets, their electronic structure modifications by the strain effect
might depress the separation of photoexcited charge carriers, and consequently,
weaken the photocatalytic performance. In addition, strain also leads to the formation
of structural defects in 2D materials. In BiOBr-square nanosheets, a large strain of
1.8 % is verified by the XRD and GPA results. It is believed that strain-induced defects
will act as recombination centres for photoexcited electrons and holes, and also depress
the quantum conversion efficiency of strained BiOBr-square nanosheets.
105
Table 5.1 BET test of the surface area of different BiOBr samples.
Sample BET (m3/g)
BiOBr-1
BiOBr-2
BiOBr-3
BiOBr-4
7.03
7.85
5.47
4.49
5.7 Summary
In summary, the strain effect on the photocatalytic activity of BiOBr nanosheets with
highly reactive {001} facets exposed was studied. The XRD, TEM, and strain tensor
simulation results reveal that the intensity and distribution of inner strain in the BiOBr
nanosheets can be modulated by adjusting the pH value in the synthesis reaction. It is
found that the strain effect can effectively tune the photocatalytic activity of BiOBr
nanosheets in dye degradation. Our work suggests that strain engineering could be an
effective approach to controlling the electronic structure of semiconductors for further
enhancement of their efficiency in converting light into other forms of energy.
5.8 References
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109
Chapter 6 Construction of 2D lateral pseudo-heterostructures by strain
engineering
6.1 Introduction
Discontinuities at interfaces in heterogeneous structures can lead to exciting and
possibly non-trivial properties, due to broken symmetries at the interfaces. Through
constructing heterostructures, one can feasibly engineer and manipulate electronic,
optical, and magnetic phases at so-called heterointerfaces, and thus, generate unusual
properties and new phenomena.[1,2] Numerous breakthroughs have been achieved by
taking advantage of vertical heterostructures in previous studies. For examples,
naturally formed heterostructures give birth to high-temperature superconductivity
(High-TC) and topological states in copper-oxide-based superconductors and
topological insulators, respectively.[3,4] Artificial heterostructures were also
constructed in laminated LaAlO3/SrTiO3 films, in which magnetic order and
superconductivity surprisingly coexist.[5,6] Very recently, fabrication of vertical van de
Waals (vdW) heterogeneous structures were successfully achieved by using atomically
thin layer materials, including graphene and h-BN, which drew immediate attention.[7-
10] Besides vertical heterostructures, lateral heterostructures with two materials joined
laterally have also drawn great attention.[11] The research in this field was boosted soon
after single-layer-thick 2D lateral heterostructures were successfully fabricated, as they
possess controllable band-offset tuning and can be used in electronics, optoelectronics,
and catalysis. Lateral interfaces in 2D lateral heterostructures are constructed from
covalent bonds and not linked by vdW forces. Less electron and phonon scattering
centers are expected across the interface. Electron hopping and band alignment in such
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lateral heterostructures are therefore less affected by the interfaces, which, in turn,
promotes charge carrier transport across the lateral interfacial junctions. These
advantages of 2D lateral heterostructures have been observed and verified in
graphene/h-BN lateral heterogeneous structures. In these systems, high field-effect
mobility of charge carriers and low hysteresis behavior have been demonstrated.[12,13]
2D lateral heterostructures based on two different transition-metal dichalcogenides
(TMD) were also successfully fabricated, exhibiting lateral p-n diodes behaviors or
staggered band alignment. A broad range of applications based on 2D lateral TMD
heterostructures, therefore, have been proposed, such as logic circuits, field-effect
transistors, and photodetectors.[14-16]
Despite their appealing advantages, great challenges remain in constructing 2D
lateral heterostructures with the desired properties. A clean and atomically sharp
interface is the essential requirement to achieve exotic characteristics. To achieve this
goal, sophisticated and costly manufacturing methods such as MBE, lithography, and
chemical vapor deposition are necessary for fabricating these high-quality
interfaces.[11,17] It is still extremely difficult to achieve 2D lateral heterostructures with
a large area, instead, most epitaxially-grown samples possess small heterogeneous
domains. The other limitation originates from lattice mismatch at the lateral
heterointerface. The candidate materials for 2D lateral heterostructures must have
similar crystal structures and lattice constants. This is the reason why very few 2D
lateral heterostructures have been reported so far.[18] An alternative strategy to
overcome these limitations is to develop a “lateral pseudo-heterostructure”.[19,20]
Generally, the pseudo-heterostructures have a single chemical component but show
spatial variation in their physical properties, which offers a promising way to build 2D
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systems that show the functionalities of 2D lateral heterostructures but overcome the
problems of lattice mismatch.
Here, we report a 2D lateral pseudo-heterostructure with an electronic heterogeneous
interface constructed from single-component BiOBr 2D nanosheets by strain
engineering. As a typical bismuth oxyhalide (BiOX, where X = Cl, Br, I),[21,22] BiOBr
has great potential for use in photoenergy conversion applications, owing to its indirect
band gap of 2.8 eV.[23-26] Taking advantage of their strain-sensitive electronic structure,
few-layer-thick BiOBr nanosheets were successfully prepared with controlled spatial
distributions of local electronic structures by manipulating the strain distribution. The
position and characteristics of the electronic heterogeneous interface can also be tuned
by adjusting the strain distribution in the nanosheets. Effective separation of charge
carriers at the electronic heterointerface is then demonstrated in this lateral pseudo-
heterostructure, owing to the appropriate band alignment at the electronic
heterogeneous interface. Its excellent photoresponse and enhanced photocurrent
suggest that such lateral pseudo-heterostructures are potential candidates for superior
optoelectronic devices.
6.2 Experimental section
To synthesize BiOBr nanosheets, 5 mmol Bismuth nitrate pentahydrate and 5 mmol
CTAB were, respectively, added into 100 mL distilled water at room temperature.
Sodium hydroxide solution (1 M) was added to the mixed solution to adjust the pH
value. The mixed solution was then stirred for 1 h, and poured into a 100 mL Teflon-
lined stainless autoclave up to 80% of the total volume. In the hydrothermal reaction,
the sealed autoclave was heated at 170 °C for 17 h, and then cooled in air. The resulting
precipitates were collected, washed with ethanol and deionized water several times,
112
and dried at 80 °C for 10 h. Bismuth nitrate pentahydrate, Sodium hydroxide, and
CTAB were purchased from Sinopharm Chemical Reagent Co., Ltd. (SCRC). All
reagents used in this work were of analytical grade and were used as received without
any further purification
The morphologies and microstructures of the as-prepared samples were characterized
by SEM (JEOL JSM-7500FA) and HAADF images were obtained by STEM (JEOL
JEM-ARM200F). A commercial AFM (Asylum Research MFP-3D) was used to
measure the morphology and surface potential of the BiOBr nanosheets by scanning
Kelvin probe microscopy (SKPM). A Pt/Ir coated n-silicon probe with resonance
frequency of 45-115 kHz and force constant of 0.5-0.95 M/m was used in the AFM
measurements. BiOBr powders were distributed on gold (50 nm) coated silicon
substrates. Before characterization of the surface potential of BiOBr samples, the
surface potential of the Pt/Ir tip was calibrated on a standard Si wafer coated with Au
film. Raman mapping investigations were carried out with a confocal Raman
spectrometer (Horiba Xplora) using a 532 nm wavelength laser as the excitation source.
The first-principles calculations were performed using the VASP code. The GGA was
applied to treat the exchange correlation energy with the PBE functional. The PAW
method was employed to describe the electron-ion interactions. A k-point sampling of
a 9 × 9 × 6 grid was generated with original Gamma-centred meshes. The cut-off energy
for the plane wave basis was 550 eV. The biaxial strain simulations were realized by
fixing the x and y axes and optimizing the z axis. Equilibrium geometries were obtained
by the minimum energy principle.
A pair of electrical contacts (Ti/Au, 5/50 nm thick) was fabricated on the a SiO2/Si
surface, by standard electron beam lithography, using a poly(methyl methacrylate)
resist and the lift-off method. BiOBr nanosheets were dispersed in ethanol and then
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dropped on the as-prepared surface followed by air-drying at 60 °C. Flakes suitable for
electrical characterization were identified by an optical microscope. Photocurrent
measurements were performed in vacuum at room temperature, using a probe station
connected to a semiconductor parameter analyser (Agilent B1500). The drain bias of
10 V and a gate bias of 80 V have been applied. The wavelength of the light source
was 450 nm. The power of the light was measured to be 50 mW·cm-2.
6.3 Structure characterization of BiOBr nanosheets
Figure 6.1 (a) AFM images of BiOBr-circle nanosheets. (b) and (c) AFM images of
BiOBr nanosheets synthesized with pH values between 7 and 2. (d) AFM images of
BiOBr-circle nanosheets. Scale bars are 1μm.
The morphologies of the BiOBr nanosheets were revealed by SEM and AFM. Figure
6.1 shows the evolution of the morphology of BiOBr nanosheet, which undergo a shape
transition from square to circle. They also indicate that BiOBr nanosheets are
approximately several micrometers in size. The concentration of hydroxide is thus
expected to be the key factor that determines the shape of the 2D BiOBr nanosheets.
Despite their different shapes, all the nanosheets display a layered structure, which
is confirmed in the cross-section TEM images in Figure 6.2. The thickness of the
BiOBr-square and BiOBr-circle show similar thickness of 15-40 nm, as shown in
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Figure 6.3. The large area-to-thickness ratio reflects their 2D nature. The interlayer
distances for the BiOBr-square and BiOBr-circle nanosheets are ~0.76 nm and ~0.78
nm, respectively, which correspond to the spacing of BiOBr (001) face. This indicates
that the BiOBr nanosheets have a preferred growth orientation normal along the [001]
direction.
Figure 6.2 (a) SEM image of BiOBr-square nanosheets. (b) and (c) Cross-sectional
TEM images of a single BiOBr-square nanosheet. (d) SEM image of BiOBr-circle
nanosheets. (e) and (f) Cross-sectional TEM of a single BiOBr-circle nanosheet.
115
Figure 6.3. (a) AFM image of BiOBr-square nanosheets. (b) The height profile
corresponding to the black line in (a). (c) AFM image of BiOBr-circle nanosheets. (d)
The height profile corresponding to the black line in (c).
Figure 6.4 XRD patterns of BiOBr-square and BiOBr-circle powders.
116
Figure 6.5 EDS mapping of elements (Bi, O, Br) distribution of BiOBr-square and
BiOBr-circle nanosheets. All elements uniformly distribute across the surface.
Our XRD, SEM, TEM, and EDS measurements verify that all BiOBr nanosheets
are pure phase without any detectable impurities or structural defects, as shown in
Figure 6.4 and Figure 6.5.
6.4 Strain induced pseudo-heterostructure
Figure 6.6 shows STEM images of BiOBr nanosheets with both square and circular
shape that were collected in HAADF mode. The atomic arrangements of the nanosheets
are clearly revealed and demonstrate their 2D single-crystal nature. Interestingly, these
BiOBr nanosheets have distinct local atomic features associated with their shapes,
especially in the areas close to the edges of the nanosheets. A homogeneous atomic
arrangement is observed across the whole BiOBr-square nanosheet, including edge
areas, as shown in Figure 6.6 (b). No obvious difference has been found in the atomic
fringes across the whole square nanosheet. The corresponding FFT pattern is shown in
Figure 6.6 (c) (pattern A across the whole nanosheet). Only one set of patterns can be
observed. In contrast, the BiOBr-circle sample exhibits an inhomogeneous atomic
arrangement across the nanosheet in Figure 6.6 (e). The areas close to the edge exhibit
relaxation in their atomic structures, although the central area retains a similar atomic
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arrangement to that in the square nanosheet. The FFT pattern, shown in Figure 6.6 (g),
exhibits a group of symmetric spots (pattern B, labelled in green) in addition to the
intrinsic spots (pattern A, labelled in yellow). These spots of pattern B with a shorter
nearest-neighbor spot distance are attributed to the area close to the edge (Area B) in
Figure 6.6 (f). Their presence reflects the considerable lattice relaxation exists in the
areas close to the edges, as illustrated in Figures 6.7.
Figure 6.6. (a) HAADF image of BiOBr-square, showing a homogeneous crystal
structure (inset: lower magnification to show whole nanosheet). (b) Enlarged image
from the selected area in (a). (c) Corresponding FFT pattern of the area in (b). (d) Strain
mapping of BiOBr-square by GPA. (e) HAADF image of BiOBr-circle, showing an
interface of the pseudo-heterostructure (inset: lower magnification to show whole
nanosheet). (f) Enlarged image from the selected area in (e). (g) Corresponding FFT
pattern of the area in (f). (h) Strain mapping of BiOBr-square by GPA. In the image of
GPA, the blue zones are under compressive strain, the red zones are under tensile strain,
and the green zones are not strained.
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Figure 6.7 Schematic diagram of the transition of the shape in (001) face exposed
BiOBr nanosheets. (a) BiBOr-square with most area of the edge of BiOBr-square are
terminated by {110} facets and the other small edge area terminated by {100} facets.
(b) BiOBr-circle nanosheet, without any specific terminated faces. Bi atoms and O
atoms are shown as purple and red balls respectively.
These relaxed areas induce a significantly inhomogeneous distribution of inner strain
in the BiOBr-circle nanosheets. It should be noted that a sharp interface can be
observed between Area A and Area B in the BiOBr-circle nanosheets, as indicated by
the dashed lines in Figure 6.6 (e) and (f). This interface can be regarded as a lateral
pseudo-heterointerface, because both areas have identical chemical component
composition but distinctly different inner strains. (see more TEM images in Figure 6.8
and Figure 6.9) In order to revealing the detailed local strain distribution, we used the
GPA method to study both square and circle nanosheets.[27,28] Figure 6.6 (d) and (h)
show GPA strain mapping results for the BiOBr-square and the BiOBr-circle
nanosheets, respectively. As shown in Figure 6.6 (d), the distribution of local strain is
uniform despite the subtle fluctuation in the BiOBr-square nanosheets. Whereas,
Figure 6.6 (h) shows that the inner strains of Area A and Area B in the BiOBr-circle
nanosheet are different, with a clear interface between these two areas (see more results
119
in Figure 6.10). Our high resolution TEM (HRTEM) images and GPA suggest that the
strain distribution in the BiOBr nanosheets can be manipulated through controlling the
concentration of hydroxide (pH value) in the hydrothermal reactions. Lateral pseudo-
heterostructures based on inhomogeneous strain distribution can be obtained in a single
BiOBr nanosheet.
Figure 6.8 (a) HRTEM images of BiOBr-circle nanosheets, showing clear interfaces
between edge area and central area. (b) HRTEM images of BiOBr-square nanosheets.
No interfaces are observed in BiOBr-square nanosheets.
Figure 6.9 (a)-(d) TEM images of BiOBr nanosheets showing the change of shape from
square to circle with the pH value change from 7 to 2. (e)-(f) HRTEM of BiOBr
nanosheets with intermediate shapes between square and circle.
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Figure 6.10 Strain distribution of BiOBr nanosheets obtained by GPA. (a) HAADF
image of BiOBr-square nanosheet. (b)-(d) Homogeneous strain distribution of BiOBr-
square. (e) HAADF image of BiOBr-circle nanosheet. (f)-(h) Inhomogeneous strain
distribution of BiOBr-circle, showing a clear interface at the pseudo-heterointerface.
Finally, based on the above experimental results, we propose a possible formation
mechanism of the lateral pseudo-heterostructure by reliving the inner strain in BiOBr-
circle nanosheets. In this synthesis method, the reactions could be described as follows
which were proposed by M. Shang et al.[29]
Bi3+ + C16H33(CH3)3N-Br → C16H33(CH3)3N-Br-Bi3+ (6.1)
C16H33(CH3)3N-Br-Bi3+ + OH- → BiOBr + C16H33(CH3)3N+ + H+ (6.2)
CTAB is proposed to act as both the template where BiOBr microplates nucleate and
grow, and the Br source. In the aqueous solution, as shown in the schematic diagram
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of reaction procedures in Figure 6.11, CTAB is lamellar structure with the surface
terminated by Br- anions. Mixed Bi3+ cations combine with CTAB at the initial
positions of the Br- anions, which will further react with OH- anions and form
multilamellar BiOBr precursors, as illustrated in Equation 6.2. Through adjusting the
concentration of OH-, the reaction speed can be controlled. A higher concentration of
OH- will result in a higher reaction speed, and vice versa. Then with the assistance of
the energy (high temperature and pressure) in the hydrothermal reaction, BiOBr crystal
could nucleate and grow on templates.
Figure 6.11 The schematic diagram of the reaction process of BiOBr nanosheets.
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In this process, BiOBr single crystal seed should take a tetragonal shape (square
shape) to maximize the expression of {001} facets and minimize the total surface
energy according to Wulff theorem, because {001} facets have the lowest surface
energy than other facets.[30-32] When the reaction speed is quick (high concentration of
OH- anions), as observed in TEM images, square nanosheets (BiOBr-square) can be
harvested, accompanying with large inner strain in BiOBr crystal. However, when the
reaction speed is slow (small concentration of OH- anions), the shape of BiOBr
nanosheets change to circle (BiOBr-circle), because in stress-free and initially defect-
free nanocrystalline materials edge relaxations and reorientations are effective ways to
release the inner strain. Therefore, even though the BiOBr nanosheets in two pH values
both have a layer structure with same crystal structure and exposed facets, but with
different growing speeds of the crystal, we could obtain BiOBr nanosheets with
different degrees of inner strain. Moreover, relieving the inner strain by relaxation near
the edge could create a lateral pseudo-heterostructure between the central and edge area
in BiOBr-circle nanosheets.
6.5 Electronic properties characterization
The local electronic structures in 2D materials can be significantly affected by local
strain.[33,34] To characterize to effect of the pseudo-heterointerface on the local
electronic properties, we carried out the KPFM measurement, which can
simultaneously acquire the surface morphologies and local work functions of BiOBr
nanosheets.
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Figure 6.12 (a) AFM image of BiOBr-square. (b) KPFM image of BiOBr-square. (c)
Corresponding line profiles of the lines in a and b, with the height as a black line and
the work function as a red line. (d) AFM image of BiOBr-circle. (e) KPFM image of
BiOBr-circle. (f) The corresponding line profiles of the lines in d and e, with the height
as a black line and work function as a red line.
As shown in Figure 6.12 (a) and (d), both square and circle BiOBr nanosheets present
a flat surface. The morphologies are consistent with TEM results. Figure 6.12 (b) and
(e) show the local work function mapping images acquired simultaneously with AFM
measurements on square and circle BiOBr nanosheets, respectively. By measuring the
CPD between the conductive tip with a constant work function and the surface, KPFM
can precisely determine the work function of materials with high resolution.[35] The
BiOBr-square nanosheet shows a homogeneous work function distribution across its
entire surface, as demonstrated by the line profiles in Figure 6.12 (c). In contrast, the
work function on the BiOBr-circle nanosheet exhibits an apparent variation between
the area near the edge and the central area, as shown in Figure 6.12 (e) and the line
profile in Figure 6.12 (f). We conjecture that the variation in the work function might
be attributable to distinct local electronic structures induced by the inhomogeneous
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distribution of local strains beside the lateral pseudo-heterointerface in BiOBr-circle
nanosheets. This was also observed in the BiOBr nanosheets with intermediate shapes,
as shown in Figure 6.13.
Figure 6.13 (a) The AFM image of the BiOBr-circle nanosheet with the shape between
square and circle. (b) The corresponding KPFM image of the BiOBr nanosheet with
only the circle area (indicated by green arrows) showing anisotropic charge
distribution. Scale bars are 1 μm.
In order to verify this hypothesis, confocal Raman spectroscopy was employed
(incident laser wavelength λ = 532 nm) to identify the correlation between the work
function and the local strain in BiOBr nanosheets, as shown in Figure 6.14. Two strong
phonon modes can be observed in the Raman spectra of the BiOBr samples. The Raman
peaks of BiOBr-square are located at 118.3 cm-1 and 160.5 cm-1, while the Raman
peaks of BiOBr-circle are located at 113.6 cm-1 and 160.5 cm-1. According to previous
reports, the peak around 113.6 cm-1 can be assigned to the A1g mode (internal Bi-Br
stretching), and the peak at 160.5 cm-1 can be assigned to the Eg mode (internal Bi-Br
stretching).[36] Although the Eg mode remains at 160.5 cm-1 for both samples, the A1g
peak of BiOBr-square possesses a blue shift of ~5 cm-1 from 113.6 cm-1 to 118.3 cm-1,
which suggests a large inner compressive strain. Raman mapping of the A1g peak on
125
both square and circle nanosheets were carried out to provide a straightforward way to
identify the strain distributions. As shown in the insets in Figure 6.14, the intensity of
the peak at 118.3 cm-1 for BiOBr-square was homogenously recorded across the whole
nanosheet, indicating the homogeneous distribution of strain. In contrast, for BiOBr-
circle, the intensity of the peak at 113.6 cm-1 is much stronger near the edge area than
in the central area.
Figure 6.14 Raman spectra of BiOBr nanosheets, with BiOBr-square shown by the
black line and BiOBr-circle by the red line. Insets are the Raman mappings at 118.3
cm-1 of BiOBr-square (top) and 113.6 cm-1 of BiOBr-circle (bottom).
This observation is supported by the fact that the relaxation of strain sprang from the
edge area to the inner central area in BiOBr-circle nanosheets. It should be notice that
during the Raman mapping, the resolution of is limited by the size of the laser spot,
126
which is around 361 nm in this work. The spatial variation in the A1g peak is on a
smaller scale than the resolution of the instrument, and so there is some averaging
between each analysis spot, which will border the Raman activate area in the
measurement. These results demonstrate that the variation of the local electronic
structure can be directly correlated with its inhomogeneous strain distribution beside
the lateral pseudo-heterointerface in the BiOBr-circle nanosheet.
6.6 DFT calculation of the strain effect on the band structure of BiOBr
We carried out DFT calculations to simulate the strain dependence of the electronic
structure in BiOBr. Figure 6.15 (a) and (b) show DFT calculation results for the pristine
and strain-applied BiOBr. When an in-plane strain (for example, the compressive strain
here) is applied, the band structure and DOS of BiOBr are significantly modulated. The
band gap is decreased by 0.2 eV in BiOBr under compressive strain (2.12 eV for
pristine BiOBr). The CBM shifts from the Z point to the M point in the irreducible
Brillouin zone, and the VBM moves from its position between the R and Z points to a
position between the Z and A points. The most attractive feature in the electronic
structure is that the band structure around the CBM is more dispersive in strained
BiOBr than in pristine BiOBr. This suggests that the excited electrons in strained
BiOBr have a small effective mass and high mobility. As a result, the total DOS of
BiOBr can also be tuned by strain, as shown in Figure 6.15 (b), in which the partial
DOS contributed by the O 2p orbitals is significantly changed. The change in the
electronic structure can be ascribed to the symmetry transition in the BiOBr crystal
under strain, which is a common phenomenon in many other semiconductors.[37] The
DFT calculations indicate that the distinct electronic properties flanking the pseudo-
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heterointerface in the BiOBr-circle nanosheets are indeed controlled by the local strain
distribution, which supports our experimental findings.
Figure 6.15 (a) Calculated band structure of BiOBr, with the pristine BiOBr shown by
the black solid line and the BiOBr with compressive strain shown by the red dashed
line. (b) Calculated DOS of BiOBr, with the pristine BiOBr shown by the solid line
and BiOBr with compressive strain shown by the dashed line.
The quantitative evaluation of strain effect on the electronic structures of BiOBr
were estimated based on the DFT calculations. As shown in Figure 6.16, the value of
Eg of BiOBr, which determines its ability of light absorption and photoresponse
threshold, varies remarkably under both compressive strain and tensile strain (vary
from 1.27 eV to 2.67 eV). Meanwhile, the effect mass of charge carries near the CBM
and VBM also exhibit strain-sensitive properties, according the result of DFT
calculations, which will greatly affect the transport properties of charge carriers. The
strain-sensitive electronic structure of BiOBr provides diversity in designing and
creating hetrostructures through strain engineering.
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Figure 6. 16 (a) The Eg values of BiOBr under different degree of strain. (b) The effect
mass of the electron at CBM and hole at VBM under different degree of strain, which
is calculated from the second order of derivative of energy with respect to wave-vector.
6.7 Performance of BiOBr nanosheets in photoelectronic devices
Figure 6.17 (a) Photocurrent of individual BiOBr nanosheets. (b) Acquiring rise and
decay time of BiOBr-square by fitting the on/off curve, inset is an image of the
nanodevice. (c) Acquiring rise and decay time of BiOBr-circle by fitting the on/off
curve, inset is the image of the device.
To verify the ability of separating photoexcited charge carriers in BiOBr nanosheets,
photodetectors were fabricated on BiOBr-square and BiOBr-circle nanosheets,
respectively, via electron-beam lithography. We measured the photocurrent and
photoresponse of both BiOBr nanosheets. As shown in Figure 6.16 (a), under visible
light irradiation (450 nm, 50 mW∙cm-2), photocurrent was generated in both BiOBr-
square and BiOBr-circle nanosheets. The photocurrent generated in the BiOBr-circle
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nanosheet is about 2 nA, which is nearly one order of magnitude larger than that of the
BiOBr-square nanosheet under the same experimental conditions.
In addition to the contrasting photocurrents in these two BiOBr nanosheets, a
difference in the photoresponse between BiOBr-square and BiOBr-circle nanosheets
can be observed. Generally, the dynamic response of the rise (Equation 6.3) and decay
(Equations 6.4 and 6.5) of photocurrent in BiOBr nanosheets can be fitted by the
following stretched exponential functions:
I = I0 - I0e(
-t
τγ)γ
(6.3)
I = I0e(
-t
τd)γ
(6.4)
I = AI0e(
-t
τd)γ
+ BI0𝑒(
−𝑡
𝜏′𝑑)𝛾′
(6.5)
where τr and τd are the relaxation time constants of the rise and decay, respectively.
The parameter γ falls into the range of 0 to 1. As γ approaches 1, the function
approaches classic single-exponential behavior without stretching. We have fitted the
photoresponse curves of the BiOBr nanosheets to reveal the dynamic responses under
visible light irradiation. As shown in Figure 6.17 (b), the τr for BiOBr-square is
calculated to be 0.24 s. The value of γ is calculated as 1, which suggests that the
generation of photocurrent in BiOBr-square under light irradiation is dominated by the
separation of photoexcited electron-hole pairs. The decay of photocurrent for BiOBr-
square can be described by a stretched exponential function (Equation 6.4), where the
value of τd is 0.22 s and the value of γ is 0.16, suggesting that the recombination of
photocurrent involves multiple energy processes, which may origin from the strain-
induced distortion of the crystal lattice. For BiOBr-circle, as shown in Figure 6.17 (c),
the rise of photocurrent could also be described using single exponential functions with
τr of 0.23 s, which is almost the same as for BiOBr-square. To describe the decay of
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photocurrent in BiOBr-circle, however, at least two components are necessary to fit the
persistent photocurrent decay, where two exponential functions (Equation 6.5) are used
to separately analyze two different photo-relaxation processes. The τd of the fast
process was determined to be 0.12 s with γ as 0.29, and the τ′d of the slow process was
calculated to be 0.72 s with γ as 0.05. During the photocurrent decay, it is well known
that photoexcited electron-hole recombination dominates the fast decay process in the
photocurrent. Hence, the photocurrent decreases very rapidly in the initial stage. The
slow process is caused by a mechanism in which the photoexcited carriers are trapped
and spatially separated by local potential fluctuations, which can suppress the electron-
hole recombination. For our samples, because both the BiOBr-square and BiOBr-circle
nanosheets exhibit defect-free characteristics, supported by the homogeneous element
distribution across the whole nanosheets. Therefore, it is believed that the slow
relaxation process in BiOBr-circle, which reflects the longer lifetime of photoexcited
charge carriers, is governed by the strain-induced pseudo-heterointerface.
The responsivity (Rλ, Equation 6.6), and the external quantum efficiency (EQE,
Equation 5.7) are two important measures for photoconductors and photodiodes.
(6.6)
(6.7)
where Iλ represents the photocurrent (Iillumination − Idark), Pλ is the light intensity, S is the
effective device area under illumination, and λ, h, c, and e are the wavelength of
incident light, the Planck constant, the speed of light, and the charge of the electron,
respectively. Based on photocurrent results, the values of Rλ are 0.28 A/W and 1.29
A/W for BiOBr-square and BiOBr-circle, respectively. The corresponding EQE are,
therefore, calculated as 0.77 × 102 % and 3.56 × 102 % (all parameters listed in Table
6.1), for BiOBr-square and BiOBr-circle, respectively.
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Table 6.1 Calculated R and EQE for BiOBr nanosheets.
Sample (nm) P (mW∙cm-2) I (nA) S (m2) R (A/W) EQE (%)
BiOBr-square 450 50 0.26 1.83 0.28 0.77 102
BiOBr-circle 450 50 1.80 2.80 1.29 3.56 102
Figure 6.18 Schematic diagram of the band alignment at the pseudo-heterointerface in
BiOBr-circle, which promotes the separation of photoexcited carriers.
We attribute the significant enhancement of the photocurrent in the BiOBr-circle
nanosheets to improved separation efficiency of photoexcited charge carriers, which is
illustrated in Figure 6.18. In the BiOBr-circle nanosheets, the inhomogeneous work
function () distribution beside the lateral pseudo-heterointerface induces band
bending across the interface.[38] Therefore, under light irradiation, photoexcited
electrons and holes will be spatially driven by this electric field and separate to two
different regions in the nanosheet, which will prolong their lifetime and consequently
132
enhance the light harvesting efficiency in the BiOBr-circle nanosheets. In contrast,
there are not any pseudo-heterointerface in BiOBr-square nanosheets. Therefore, we
do not expect the same scenario occurring in square nanosheets.
6.8 Summary
To summarize, the lateral pseudo-heterostructure on individual BiOBr nanosheets
was fabricated by strain engineering. The local strain distribution has been revealed by
TEM and Raman mapping. The dependence of local strain on electronic structures has
been revealed by DFT calculations. An inhomogeneous work function distribution is
observed across the lateral pseudo-heterostructure, which favors the separation of
photoexcited charge carriers under visible light irradiation. Enhanced photocurrent (by
one order of magnitude) has been achieved in BiOBr pseudo-heterostructure-based
photodetectors, which is attributed to the appropriate band alignment across the
pseudo-heterointerface. The growth of such lateral pseudo-heterostructures, thus,
offers a promising way to build systems that show the functionalities of 2D lateral
heterostructures but overcome the problems of lattice mismatch.
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Chapter 7 Summary and outlooks
OV defects were carefully introduced into rutile(110) single crystal by annealing in
vacuum conditions in the temperature range from 900 K to 1300 K, which were
checked by in-situ LT-STM. Isolated OV point defects were found on the rutile(110)
(1×1) surfaces, which are favorable to be occupied by water molecule or OH. At 1300
K, the surface experienced the evolution from (1×2) to (1×2) surfaces, which were
dominated by cross-links defects.
In further work, we demonstrate that inactive pure TiO2 rutile(110) single crystal can
be activated towards the HER by creating OV point defects. OVs and Ti3+ ions in the
surface region dominate the electrical conductivity of reduced TiO2 and the amount of
electrocatalytic active sites. Combining the well-characterized atomic surface structure
and theoretical calculations, we conclude that subOVs can promote the electron transfer
and hydrogen desorption in the electrocatalytic HER on reduced TiO2 in alkaline
media. It helps to elucidate the fundamental mechanism of the electrocatalytic activity
of reduced oxides towards the HER, which is of immense fundamental and practical
importance towards an in-depth understanding and rational optimization of TMOs as
electrocatalysts in alkaline media. Considering that all the electrocatalytic
characterization of TiO2 were measured using the single crystals, which could benefit
in controlling defect species and studying their roles in electrocatalytic reactions,
synthesis TiO2 materials in nano-scale will largely enhance the number of surface
active sites and improve their performances. In addition, TiO2(110) surface with (1×2)
reconstruction has a complicated surface structure and cross-links defects, which could
alter the surface electrocatalytic properties and electrical conductivity, providing an
opportunity to further tune and enhance the electrocatalytic activities of TiO2.
138
Another focus was by the mean of precisely tunning the growth conditions, the strain
can be controllably introduced into 2D BiOBr nanosheets with highly reactive {001}
facets exposed. It is found that the strain effect can effectively tune the photocatalytic
activity of BiOBr nanosheets in dye degradation. Our work suggests that strain
engineering could be an effective approach to controlling the electronic structure of
semiconductors for further enhancement of their efficiency in converting light into
other forms of energy.
By characterizing single nanosheets with STEM, lateral pseudo-heterostructure was
confirmed in BiOBr-circle nanosheets. The local strain distribution has been revealed
by STEM and GPA analysis. The dependence of local strain on electronic structures
has been revealed by DFT calculations. An inhomogeneous work function distribution
is observed across the lateral pseudo-heterostructure, which favors the separation of
photoexcited charge carriers under visible light irradiation. Enhanced photocurrent (by
one order of magnitude) has been achieved in BiOBr pseudo-heterostructure-based
photodetectors, which is attributed to the appropriate band alignment across the
pseudo-heterointerface. The growth of such lateral pseudo-heterostructures, thus,
offers a promising way to build systems that show the functionalities of 2D lateral
heterostructures but overcome the problems of lattice mismatch.
139
Appendix A: List of publications
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Nano 2018, 12, 5059.
(2) H. Feng, Z. Xu, L. Ren, C. Liu, J. Zhuang, Z. Hu, X. Xu, J. Chen, J. Wang, W.
Hao, Y. Du, S. X. Dou, Activating titania for efficient electrocatalysis by vacancy
engineering, ACS Catal. 2018, 8, 4288.
(3) J. Zhuang, C. Liu, Z. Zhou, G. Casillas, H. Feng, X. Xu, J. Wang, W. Hao, X.
Wang, S. X. Dou, Z. Hu, Y. Du, Dirac signature in germanene on semiconducting
substrate, Adv. Sci. 2018, 1800207.
(4) Z. Xu, K. Xu, H. Feng, Y. Du, W. Hao, Sp orbital hybridization: a strategy for
developing efficient photocatalysts with high carrier mobility, Sci. Bull. 2018, 63, 465.
(5) P. De Padova, H. Feng, J. Zhuang, Z. Li, A. Generosi, B. Paci, C. Ottaviani, C.
Quaresima, B. Olivieri, M. Krawiec, Y. Du, Synthesis of multilayer silicene on
Si(111)√3 × √3-Ag, J. Phys. Chem. C 2017, 121, 27182.
(6) H. Feng, J. Zhuang, A. D. Slattery, L. Wang, Z. Xu, X. Xu, D. Mitchell, T. Zheng,
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dimensional lateral pseudoheterostructures by strain engineering, 2D Mater. 2017, 4,
025102.
(7) H. Feng, Y. Du, C. Wang, W. Hao, Efficient visible-light photocatalysts by
constructing dispersive energy band with nisotropic p and s-p hybridization states,
Curr. Opin. Green Sustain. Chem. 2017, 6, 93.
140
(8) Y. Liu, H. Feng, X. Yan, J. Wang, H. Yang, Y. Du, W. Hao, The origin of enhanced
photocatalytic activities of hydrogenated TiO2 nanoparticles, Dalton Trans. 2017, 46,
10694.
(9) X. Lou, J. Shang, L. Wang, H. Feng, W. Hao, T. Wang, Y. Du, Enhanced
photocatalytic activity of Bi24O31Br10: constructing heterojunction with BiOI, J. Mater.
Sci. Technol. 2017, 33, 281.
(10) Z. Li, H. Feng, J. Zhuang, N. Pu, L. Wang, X. Xu, W. Hao, Y. Du, Metal-silicene
interaction studied by scanning tunneling microscopy, J. Phys.: Condens. Matter 2016,
28, 034002.
(11) H. Feng, X. Xu, W. Hao, Y. Du, D. Tian, L. Jiang, Magnetic field actuated
manipulation and transfer of oil droplets on a stable underwater superoleophobic
surface, Phys. Chem. Chem. Phys. 2016, 18, 16202.
(12) H. Liu, H. Feng, Y. Du, J. Chen, K. Wu, J. Zhao, Point defects in epitaxial silicene
on Ag(111) surface, 2D Mater. 2016, 3, 025034.
(13) Y. Du, J. Zhuang, J. Wang, Z. Li, H. Liu, J. Zhao, X. Xu, H. Feng, L. Chen, K.
Wu, X. Wang, S. X. Dou, Quasi-freestanding epitaxial silicene on Ag(111) by oxygen
intercalation, Sci. Adv. 2016, 2, e1600067.
(14) L. Ren, J. Zhuang, G. Casillas, H. Feng, Y. Liu, X. Xu, Y. Liu, J. Chen, Y. Du, L.
Jiang, S. X. Dou, Nanodroplets for stretchable superconducting circuits, Adv. Funct.
Mater. 2016, 26, 8111.
(15) R. J. Wong, J. Scott, G. K. C. Low, H. Feng, Y. Du, J. N. Hart, R. Amal,
Investigating the effect of UV light pre-treatment on the oxygen activation capacity of
Au/TiO2, Catal. Sci. Technol. 2016, 6, 8188.
141
(16) Z. Xu, W. Hao, Q. Zhang, Z. Fu, H. Feng, Y. Du, S. Dou, Indirect-direct band
transformation of few-layer BiOCl under biaxial strain, J. Phys. Chem. C 2016, 120,
8589.
(17) H. Feng, Z. Xu, L. Wang, Y. Yu, D. Mitchell, D. Cui, X. Xu, J. Shi, T. Sannomiya,
Y. Du, W. Hao, S. Dou, Modulation of photocatalytic properties by strain in 2D BiOBr
nanosheets, ACS Appl. Mater. Interfaces 2015, 7, 27592.
(18) J. Zhuang, X. Xu, H. Feng, Z. Li, X. Wang, Y. Du, Honeycomb silicon: A review
of silicene, Sci. Bull. 2015, 60, 1551.
(19) X. Xu, J. Zhuang, Y. Du, H. Feng, N. Zhang, C. Liu, T. Lei, J. Wang, M. Spencer,
T. Morishita, X. Wang, SX. Dou, Effects of oxygen adsorption on the surface state of
epitaxial silicene on Ag(111), Sci. Rep. 2014, 4, 7543.
Appendix B: Conference contributions
1. 2018 international symposium on advanced materials and sustainable technology,
Brisbane, Oral presentation: H. Feng, W. Hao, Y. Du, S. X. Dou, Surface defect
engineering in semiconducting (photo)electrocatalyst, 2018.
2. The international conference on nanoscience and nanotechnology, Wollongong,
Poster section: H. Feng, Z. Xu, Z. Hu, W. Hao, S. X. Dou, Y. Du, Vacuum reduced
TiO2 for electrochemical HER in alkaline condition, 2018.
3. The 3rd international conference on 2D materials and technology, Singapore, Poster
section: H. Feng, L. Wang, Z. Xu, J. Zhuang, X. Xu, Y. Du, W. Hao, S. X. Dou,
Construction of 2D lateral pseudoheterostructures by strain engineering, 2017.
ICON-2DMAT Best poster award.
4. The 2nd international symposium on renewable energy technology, Sydney. Poster
section: H. Feng, Z. Xu, L. Wang, W. Hao, S. X. Dou, Y. Du, Modulation of
photocatalytic properties by strain in 2D BiOBr nanosheets, 2016. Best poster reward.