List of symbols and abbreviations - Shodhgangashodhganga.inflibnet.ac.in/bitstream/10603/90094/15/15_list_of_publications.pdf2.2.2 Applications of SiC nanostructures 59 Bibliography

ARC PLASMA SYNTHESIS OF SILICON AND

SILICON CARBIDE NANOSTRUCTURES:

CHARACTERIZATION AND APPLICATIONS

A THESIS SUBMITTED TO THE

SAVITRIBAI PHULE PUNE UNIVERSITY

FOR THE DEGREE OF

DOCTOR OF PHILOSOPHY

IN

PHYSICS

BY

Miss. CHITI MANOHAR TANK

UNDER THE GUIDANCE OF

Dr. V. L. MATHE

DEPARTMENT OF PHYSICS


PUNE 411007, INDIA

Prof. (Mrs) S. V. BHORASKAR

EMERITUS SCIENTIST

DEPARTMENT OF PHYSICS


PUNE 411007, INDIA

MARCH 2015

DETAILS OF THE PH. D. PROGRAM

Title of the Thesis ARC PLASMA SYNTHESIS OF SILICON AND

SILICON CARBIDE NANOSTRUCTURES:

CHARACTERIZATION AND APPLICATIONS

Name of the Candidate Miss. Chiti Manohar Tank

Name of the Research Supervisor Dr. V. L. Mathe

Name of the Research Co-Supervisor Prof. (Mrs.) S. V. Bhoraskar

Research Program

CSIR Emeritus Scientist Scheme and CSIR Direct

SRF

Date of Registration 20th

May 2010

CERTIFICATE

It is certified that the work incorporated in this thesis entitled “Arc Plasma Synthesis

of Silicon and Silicon carbide Nanostructures: Characterization and Applications” submitted

by Miss. Chiti Manohar Tank was carried out by the candidate under our supervision. The

work incorporated in this thesis has not been submitted to any other University or Institute

for the degree of Ph. D or any other degree or academic award. Such materials, as has been

obtained from other sources, have been duly acknowledged in the thesis.

Dr. (Mrs.) S.V. Bhoraskar

(Research Supervisor)

Emeritus Professor,

Department of Physics,

Savitribai Phule Pune University,

Pune - 411007, INDIA

([email protected] &

[email protected])

Dr. V. L. Mathe

(ResearchGuide)

Assistant Professor,



Pune 411007,INDIA

([email protected] ,

[email protected] )

Place……………

Date…………….

mailto:[email protected]




DECLARATION

I hereby declare that the present thesis entitled “Arc Plasma Synthesis of Silicon and

Silicon Carbide Nanostructures: Characterization and Applications” is an account of original

work carried out by me. This work or part(s) of the work thereof has not been submitted to

any other University or Institute for the award of any degree or diploma.

Miss. Chiti Manohar Tank (Candidate)

Senior Research Fellow,



Pune 411007,

INDIA,

([email protected])

Forwarded through,

Dr. (Mrs.) S.V. Bhoraskar

(Research Supervisor)

Dr. V. L. Mathe

(ResearchGuide)

Place………….

Date…………..


Acknowledgement

While starting any new expedition in life the most

important thing we require is able guidance and support, to

reach the proper destination. In my journey of research this place

was taken by my guide “Dr. V. L. Mathe” and co-guide “Prof. S. V.

Bhoraskar”. Dr. V. L. Mathe is a noble person who always kept on

encouraging and providing all kind of support in lab for good

work and taking care of avoiding any kind of shortage which

would create obstacle while working in the lab. Prof. S. V.

Bhoraskar is a very sweet person ‘like young forever’ with great

enthusiasm, scientific temper with relentless energy for teaching,

guiding and helping students. I was greatly benefitted by the

scientific discussions with both of them. Both of them always work

hard to maintain a good temperament for research in lab. It

would have been very difficult for me to learn different scientific

techniques and to know the world of research in their absence. I

heartily thank them for guiding me through this journey and

ask for their blessings. And I would always miss them and the

discussions with them.

After the guide, come the senior and junior members of the

lab who make the environment of the lab workable. I was lucky

that I got all the good lab mates who in some or the other aspect

helped me. They were Naveen Kulkarni and Ashok Nawale who

taught me to operate the plasma systems. It was Vijay Varma who

taught me use of many softwares and provided information

about using them. Nilesh, Vijay and Suyog assisted in some of

experiments and we had fruitful scientific discussions. They were

Supriya, Suyog, Nilesh and Gayatri with whom I shared some

wonderful moments. All the other labmates were also very helpful.

During this journey of research I also had under me some very

good project students, teaching them and working with them was

a good time few to name are Swapnil and Deepali. Thanks to all

of them for everything.

The place Department of Physics, SP Pune University is one of

the most wonderful place to work. The good culture that is

followed here is there is no discrimination made between

students and faculty. This reduces the gap between them and

gives a free environment for every student to share and discuss

Science with anyone. Everyone here is always ready for a fair

discussion about science and advancements in science. I am very

obliged that I got an opportunity to be a part of this place and

thank all the faculties. Also, I am thankful for the facilities

department has provided. Also, I acknowledge previous and

current heads of departments Prof. P. B. Vidyasagar and Prof. S.

I. Patil who always were helpful and always took note of different

characterization facilities and tried to involve more new

facilities to the department. Apart from teaching staff,

nonteaching staff was also very helpful in every aspect. I thank

all of them Mrs Dikshit, Mrs Shiekh, Mr Padvi, Mrs Kalpana, Mr

Ghule, Mr Lolage, Mr Jagtap, Mr. More, Mr Kadam, Mr Bhujbad

and all others for their timely help and support.

I would give special thanks to Lolage Sir who maintained

UV-Visible Spectrometer, FTIR and XRD and helped in recording

data. Also, I would acknowledge Shridhar Krishna who recorded

the TEM micrographs as much as I wanted and provided the TEM

software for analysis. Pore Sir was always there for any help

regarding software and hardware of computer, heartily thanks

to him for that!!

I would also add that it was university workshop where

timely support was received for small machining works. This was

a great help from all in the workshop.

It was a great help of Dr. N. P. Lalla in carrying out some

TEM measurements at UGC-DAE Consortium for Scientific

Research, Indore. Also I am thankfull to Dr. V. Sathe UGC-DAE

Consortium for Scientific Research, Indore, for providing Raman

Spectroscopy data. Scanning Tunneling Microscope (STM)

measurements were carried out at Institut für Experimentelle

und Angewandte Physik, Christian-Albrechts-Universität zu Kiel,

D-24098 Kiel, Germany by Sujoy Karan and analyzed with the

help of Sujoy Karan and Prof. Richard Berndt. Energy Filtered-

High Resolution Transmission Electron Microscopy (EF-HRTEM)

and Nanobeam Electron Spectroscopy (NES) and Diffraction

(NED) measurements were carried out and analyzed at

Dipartimento di Fisica, Università Roma Tor Vergata and Unità

CNISM, via della Ricerca Scientifica 1, 00133 Roma, Italy and

Dipartimento di Tecnologie e Salute, Istituto Superiore di Sanità,

00161 Roma, Italy with the help of Dr. Paola Castrucci, Dr. Marco

Diociaiuti, Stefano Casciardi, Francesca Tombolini, Manuela

Scarselli and Prof. Maurizio De Crescenzi. Dr. Sujatha Raman

and Prof. S. W. Gosavi helped and guided in carrying out

antimicrobial studies while Padmashree Joshi and Joag Sir

helped in carrying out field emission studies. I heartily thank all

of them for their scientific help. It was good time working with

Sujatha maam and Padmashree.

Since we are born, our parents and family members are

with us for all sort of support. So, thanking them in words is very

small. They were my parents who taught me to learn and

encouraged me in whatever I wanted to learn and my sisters,

Darsha and tvisha who always stood by my side. It was Sujoy who

taught me many minor things and kept on motivating me.

So, this is how my Ph. D. work completed with inputs from so

many people and I owe true gratitude for all of them.

Chiti Tank

A true researcher never says the work is ultimate, but always sees scope

for further improvement!!!!

Contents

Chapter 1 1

Introduction and Scientific Background 1

1.1 Introduction 2

1.2 Objective of thesis 6

1.3 Organization of thesis 7

1.4 Scientific background 7

A) Synthesis methods 7

1.4.1 Arc plasma 7

1.4.2 Mechanism of thermal plasma assisted synthesis of nanoparticles 14

B) Materials 18

1.4.3 Silicon 18

1.4.4 Silicon carbide 31

Bibliography 38

Chapter 2 43

Literature Survey 43

2.1 Silicon 44

2.1.1 Synthesis methods and applications of silicon nanoparticles (SiNPs) 45

2.1.2 Synthesis methods and applications of silicon nanowires (SiNWs) 47

2.1.3 Synthesis Methods and Applications of Silicon Nanotubes (SiNTs) 50

2.2 Silicon carbide 54

2.2.1 Synthesis of SiC nanostructures 55

2.2.2 Applications of SiC nanostructures 59

Bibliography 60

Chapter 3 70

Experimental Techniques & Procedures 70

3.1 Experimental method of synthesis 71

3.1.1 DC direct arc thermal plasma set up 71

3.1.2 Synthesis procedure and mechanism of synthesis 75

3.2 Characterization techniques 76

3.2.1 Transmission electron microscopy 78

3.2.2 Electron energy-loss spectrometry in TEM 86

Bibliography 90

Chapter 4 91

4.1 Synthesis and characterization of silicon nanotubes 92

4.1.1 Experimental details 92

4.1.2 Results and discussion 94

A) Samples synthesized in presence of argon 94

B) Samples synthesized in presence of argon and hydrogen (95:5 mole%) 96

4.1.3 Conclusions 108

4.2 Synthesis of silicon nanostructures in presence of different hydrogen

concentrations and its effect on the morphology 108




4.3 Antibacterial study of silicon nanoparticles (Si1) and nanotubes (Si5) 115

4.3.1 Introduction 115




4.4 Field emission study of silicon nanotubes (Si5) 122

4.4.1 Introduction 122

4.4.2 Electron field emission 122

4.4.3 Experimental procedure for field emission study 123



Bibliography 126

Chapter 5 130

Synthesis of Silicon carbide Nanostructures & Application 130

5.1 Introduction 131

5.2 Synthesis and characterization of SiC nanoparticles 132




5.3 SiCNPs - diglycidyl ether bisphenol A (DGEBA) epoxy polymer composites 153

5.3.1 Epoxy polymers 153

5.3.2 Diglycidyl ether bisphenol A (DGEBA) 153

5.3.3 Curing of epoxy 154

5.3.4 Procedure of preparation of SiC nanoparticles - DGEBA composites 155

5.3.5 Study of properties of SiCNP – DGEBA composites 157


Bibliography 162

Chapter 6 165

6.1 Conclusions 166

6.2 Future Scope 168

List of symbols and abbreviations


Symbols/

Abbreviations

Full form/ Meaning

NPs Nanoparticles

NTs Nanotubes

NSs Nanostructures

NWs Nanowires

SiNPs Silicon Nanoparticles

SiNSs Silicon Nanostructures

SiNTs Silicon Nanatubes

SiNWs Silicon Nanowires

SiCNPs Silicon Carbide Nanoparticles

SiCNSs Silicon Carbide Nanostructures

SiCNTs Silicon Carbide Nanatubes

SiCNWs Silicon Carbide Nanowires

CNTs Carbon Nanotubes

DGEBA Diglycidyl Ether bisphenol A

β-SiC or 3C-SiC FCC diamond SiC

α-SiC Hexagonal polytypes of SiC

LO Longitudinal optical phonon

TO Transverse optical phonon

QC Quantum confinement

PL Photoluminescence

PS Porous Silicon

CVD Chemical Vapour Deposition

DC Direct Current

AAO Anodized Aluminium Oxide

CF Conflat Flange


KF Kwik flange

XRD X-Ray Diffraction

UV-Vis UV-Visible

FTIR Fourier Transform Infrared

SEM Scanning Electron Microscopy

TEM Transmission Electron Microscopy

HRTEM High Resolution Transmission Electron Microscopy

EF-HRTEM Energy Filtered – High Resolution Transmission Electron Microscopy

EELS Electron Energy – Loss Spectroscopy

NEELs Nano-beam Electron Energy Loss Spectroscopy

SAED Selective Area Diffraction Pattern

NED Nano-beam Electron Diffraction

TG Thermogravimetry

TGA Thermogravimetry Analysis

RF Radio Frequency

ICP Inductively Coupled Plasma

VLS Vapour Liquid Solid

CFU Colony Forming Units

UHV Ultra-high Vacuum

FE Electron Field Emission

HOPG Highly oriented pyrolytic graphite

List of Figures

List of Figures

1.1 (a) Plot of number of publications per year obtained from ‘web of science’ on

searching ‘nano’ (b) Public R & D investments in nanotechnology globally [1]. 2

1.2 Plot of number of publications on silicon published per year (searched word silicon

in Science Direct). 4

1.3 Plot of publications and scitations (cumulative) with year by searching for ‘thermal

plasma’ (http://academic.research.microsoft.com/). 6

1.4 (a) Current-voltage characteristic of a DC discharge through a gas [22], and (b) the

dependence of individual specie temperature in thermal plasma on pressure. 8

1.5 Schematic of the characteristic arc-regions in an unspecified electric-arc. 11

1.6 Schematic distribution of arc-voltage along an unspecified arc-length. 11

1.7 (a) Formation of the embryo: G1 and G2 are the Gibbs Free energies before and after

the formation of embryo, (b) the Free energy change associated with homogeneous

nucleation with radius r at different temperatures, (c) addition of atoms from the

parent phase into the interface of a critical nucleus, and (d) the temperature

dependence of nucleation rate I and growth rate U [28]. 15

1.8 sp3 hybridization in silicon. 19

1.9 (a) FCC diamond crystal structure, and (b) the first Brillouin zone of the FCC lattice

with points of symmetry shown. 19

1.10 Schematic of the energy bands of silicon in the energy range near the forbidden

energy gap at 300K [31]. 20

1.11 (a) Schematic of direct indirect optical transitions in Si [32], and (b) absorption

spectra of single-crystal silicon at 77 K and 300 K [33]. 20

1.12 (a) Schematic of graphene sheet and carbon nanotubes, and (b) schematic of

stacking observed in Si due to sp3 hybridization. 26

1.13 The figures of zig-zag (left) and armchair (right) silicon nanotubes structure [55]. 27

1.14 (a) Square, pentagonal, and hexagonal single-walled SiNTs proposed by Bai et al.

[19], (b) the antiprismatic, prismatic and chiral SiNTs proposed by Lee et al. [60] 29

1.15 (a) An irregular quadrilateral lattice [61], and (b) side and front views of optimized

structures of infinite SiNTs clean (top), puckered with hydrogen capped on inner

and outer surfaces (middle), and ), with hydrogen capped on outer surface [62]. 29

List of Figures

1.16 Cluster-assembled hydrogen passivated SiNTs proposed by Guo et al. [64] 30

1.17 Primitive hexagonal unit cells of the most simple SiC polytypes. Si atoms are

represented by open circles, C atoms by filled circles. Bilayers of the three possible

positions in projection with the c-axis are labeled by the letters A, B, and C. The Si-

C bonds in the (11 0) plane indicating the relative shifts of the bilayers are

represented by heavy solid lines. The figure has been reproduced from K¨ackell et

al. (1994)[67]. 32

3.1 (a) The schematic of the DC Direct arc plasma reactor used for the synthesis of

silicon and silicon carbide nanostructures (b) the photograph of the DC Direct arc

plasma reactor. 72

3.2 The schematic of electrode assembly used for synthesis of silicon nanoparticles (a)

Anode assembly showing SS hollow rod, copper cup and cylindrical graphite

crucible (CR1) marked by 1, 2 and 3 respectively and (b) cathode consisting of

tungsten rod marked by 4. 73

3.3 The photographs of the electrode assembly used for synthesis of silicon

nanoparticles (a) Anode assembly showing SS hollow rod, copper cup and

cylindrical graphite crucible (CR2) marked by 1, 2 and 3 respectively and (b)

cathode consisting of tungsten rod marked by 4. 73

3.4 The schematic of electrode assembly used for synthesis of silicon carbide

nanoparticles (a) Anode assembly showing SS hollow rod, copper cup and conical

graphite crucible marked by 1, 2 and 3 respectively, (b) Anode assembly showing

SS hollow rod, copper cup and first stage graphite crucible marked by 1, 2 and 3

respectively, 4a (CR3), 4b (CR4) and 4c (CR5) represent second stages of crucibles

and (b) cathode consisting of tungsten rod fitted with a graphite cap marked by 4. 74

3.5 The photograph of electrode assembly used for synthesis of silicon carbide

nanoparticles (a) Anode assembly showing SS hollow rod, copper cup and conical

graphite crucible marked by 1, 2 and 3 respectively, (b), (c) and (d) consist of anode

assembly showing SS hollow rod, copper cup and two stage graphite crucibles

marked by 3 (first stage) and 4a(CR3), 4b (CR4) and 4c (CR5) and (e) cathode

consisting of tungsten rod fitted with a graphite cap marked by 4. 75

3.6 The schematic of the process of growth induced by thermal plasma [6].

76

List of Figures

3.7 Schematic representation of contrast generation depending on the mass and the

thickness of a certain area [10]. 79

3.8 (a) Left: bright-field mode, and (b) Right: dark-field mode [11]. 81

3.9 (a) Left: ray diagram to obtain selective area diffraction pattern in TEM, (b) Right:

geometry for electron diffraction and definition of camera-length, L. The electron

wavelength is λ, and the camera constant of (eqn 3.6) is λL [11]. 82

3.10 Diffraction pattern for FCC Si crystal obtained using software Carine

Crystallography 3.1 oriented in different directions. (a) (100) Zone axis, (b) (101)

Zone axis, (c) (111) Zone axis, (d) (211) Zone axis (e) (311) Zone axis and (f) (331)

Zone axis. 84

3.11 Diffraction pattern for hexagonal lattice with ABAB stacking sequence obtained

using software Carine Crystallography 3.1 oriented in different directions, (a) (001)

Zone axis, (b) (101) Zone axis, (c) (110) Zone axis, (d) (100) Zone axis 85

3.12 Energy-loss spectrum of an iron fluoride film: (a) low-loss region with a

logarithmic intensity scale and (b) part of the core-loss region, with linear vertical

scale [13]. 87

3.13 Energy-loss spectra recorded from silicon specimens of two different thicknesses.

The thin sample gives a strong zero-loss peak and a weak first-plasmon peak; the

thicker sample provides plural scattering peaks at multiples of the plasmon energy

[12]. 89

4.1 X-Ray diffraction pattern of Si samples synthesized in ambient argon. 95

4.2 TEM micrographs of as synthesized Si samples in argon (a) Si1, (b) Si2, (c) Si3 and

(d) Si4 (Insets show the selective area electron diffraction pattern of the

corresponding samples). 96

4.3 X-Ray diffraction pattern of as synthesized Si samples in presence of argon and

hydrogen in the ratio (95:5). 97

4.4 TEM micrograph of sample Si5, right inset shows the magnified image of a

nanotubes and left inset shows the corresponding SAED pattern of the nanotubes

and nanoparticles. 98

4.5 (a) TEM micrograph of sample Si6 where the inset shows the corresponding SAED

pattern, and (b) the magnified image of the tip of an elongated structure.

99

List of Figures

4.6 TEM micrograph of sample Si7 and the inset shows the corresponding SAED

pattern. 99

4.7 (a) and (b) TEM micrographs of sample Si8 and the inset in (a) shows the

corresponding SAED pattern. 100

4.8 (a) STM image (360 nm X 360 nm) of single silicon nanotube on HOPG; Vbias = 1

V, Itunn= 0.95 nA, (b) line profile along the yellow line drawn in (a). 100

4.9 Raman Spectra of silicon samples (a) crystalline silicon, (b) sample Si5. 101

4.10 Nanobeam low electron energy loss spectra for two different nanotubes (curves (a)

and (b)), a spherical nanoparticle (curve (c)) and a SiO2 standard (curve (d)); inset:

the complete experimental SiO2 NEELS spectrum presenting the zero-loss peak due

to elastically transmitted electrons and first order plasmon features at energies

between 10 and 30 eV. 102

4.11 Si L2,3 edge electron energy – loss spectra recorded for the SiO2 specimen (curve

(a)), a spherical nanoparticle (curve (b)), two nanotubes (curves (c) and (d)), and the

clean Si nanotube (curve (e)). 103

4.12 EF-HRTEM image of a nanotube. The upper left inset reports the FFT of the area

contained in the white square; the lower right inset shows the filtered image of the

region obtained by making the inverse of the FFT displayed in the upper left inset. 104

4.13 NED of the nanotube imaged in the upper left inset of the .The bright circular area

indicates the region from which diffraction pattern arises. In the upper right inset the

profile of the diffraction pattern obtained along a straight line passing through its

center is reported. 106

4.14 (a) EF-HRTEM image of a nanoparticle; the inset shows the FFT calculated for the

white square region. (b) NED of the same nanoparticle also imaged in the upper left

inset of the figure. The bright circular area indicates the region from which

diffraction pattern arises. In the upper right inset the profile of the diffraction pattern

obtained along a straight line passing through its center is reported. 107

4.15 X-Ray diffraction patterns of samples synthesized in increasing H2 – concentration. 110

4.16 TEM micrograph of silicon nanostructures (a) S1, (b) S2, (c) S3, and (d) S4. 111

4.17 TEM images of silicon nanowires. (a) and (b) lattice spacing on silicon nanowires

showing twin boundary, (c) lattice spacing on silicon nanowires and (d) TEM image

of the mouth of a nanowire showing lattice spacing of 1.94 Å. 112

List of Figures

4.18 (a) TEM image of spherical nanoparticles of silicon observed in Sample S3, (b)

magnified image showing lattice planes and (c) fast Fourier transform of image (b). 113

4.19 HRTEM image of hexagonal platelet of silicon carbide; lower left inset shows the

magnified image of the region marked by square and lower right inset show the

corresponding fast Fourrier transform. 113

4.20 FTIR spectra of silicon nanostructures synthesized in different gas compositions. 114

4.21 Effect of different concentrations of silicon nano-structures on bacterial strains

tested. Standard absorbance values of silicon nano-structures at various

concentrations from 0 µg/ml (positive control) to 200 µg/ml are provided.

Experimental mixture having NB media with respective bacterial inoculums,

without nano-structures was used as positive control. NB media alone was used as

negative control (a) Effect of Si1 on bacterial strains tested (b) Effect of Si5 on

bacterial strains tested. 119

4.22 Colony forming units counting in Gram-positive bacterial strains calculated for

different concentrations of nano-structures (0 to 200 µg/ml) (a) CFU of B. subtilis

cultures calculated at the dilutions of 103, 10

4 and 10

5 for Si1, (b) CFU of B. subtilis


4 and 10

5 for Si5, (c) CFU of S. aureus


8 and 10

9 for Si1 and (d) CFU of S.

aureus cultures calculated at the dilutions of 107, 10

8 and 10

9 for Si5. 120

4.23 Colony forming units counting in Gram-negative bacterial strains calculated at the

dilutions of 107, 10

8 and 10

9 for different concentrations of nano-structures (0 to 200

µg/ml) (a) CFUs of E-coli cultures for Si1, (b) CFU of E-coli cultures for Si5, (c)

CFU of P. aeruginosa cultures for Si1 and (d) CFU of P. aeruginosa cultures for

Si5. 121

4.24 (a) SEM image of SiNT coated W- tip, (b) J-E plot (inset shows FN plot), (c)

emission current vs. time plot and (d) FEM micrograph. 124

5.1 (a) Plot of rate of change in weight of anode (Awl) and cathode (Cwl) for samples

synthesized using different crucible shapes at 80 A arc current, (b) Plot of Awl and

Cwl for samples synthesized using different crucible shapes at 100 A arc current, (c)

Plot of yield (Y) and ratio of Y to Awl (β) for samples synthesized using different

crucible shapes at 80 A arc current, and (d) Plot of Y and β for samples synthesized

using different crucible shapes at 100 A arc current. 135

List of Figures

5.2 X-Ray diffraction patterns of SiC- nanoparticle samples synthesized by thermal

plasma. 137

5.3 Thermogravimetric graphs of all as synthesized SiC samples. 141

5.4 TEM micrographs of samples (a) SiC1 and (b) SiC2. 143

5.5 TEM micrographs of samples (a) SiC3 and (b) SiC4. 143

5.6 TEM micrographs of (a) carbon hollow and graphene like structures and (b)

graphitic nanostructures. 144

5.7 TEM micrographs of (a) as synthesized sample SiC7, (b) as synthesized sample

SiC9, (c) heat treated sample SiC7 and (d) heat treated sample SiC9. (Insets show

the SAED patters of the corresponding samples). 145

5.8 (a) TEM micrograph of SiC sample showing typical faceted structures (b) HRTEM

image of the hexagonal 2D structure which is further magnified in (c) with its FFT

image in (d), (e) TEM micrograph of single hexagonal structure with corresponding

SAED pattern in inset, (f) Schematic showing possible growth direction resulting in

the formation of hexagonal 2D structure, (g) Schematic showing possible growth

direction resulting in the formation of triangular 2D structure, (h) Schematic

showing possible growth direction resulting in the formation of triangular pyramidal

structure. 146

5.9 TEM micrograph showing different structures of SiC nanoparticles. Insets 1, 2, 3

and 4 show FFT from the region marked by square 1, 2, 3 and 4. 148

5.10 (a) TEM micrograph of triangular shaped SiC nanoparticles, (b) TEM micrograph

of same triangular shaped SiC nanoparticles from different view, (c) TEM

micrograph of a structure observed in SiC samples, (Upper insets show HRTEM

images of red squares and lower insets show the corresponding FFT image). 148

5.11 (a) TEM micrograph of SiC-Si nanojunction formation (lower hexagonal sheet

belongs to SiC while the hemispherical structure belongs to Si), (b1) and (b3) show

the magnified images of square 1 and 2 in (b) and (b2) and (b4) show the

corresponding FFT images. (c) TEM micrograph consisting of Si and SiC junction,

(c1) FFT of square 1in (c) showing presence hexagonal Si, and (c2) FFT of square

2 in (c) showing presence of hexagonal silicon carbide. 149

5.12 UV-Visible absorption spectra of samples SiC1 and SiC2.

150

List of Figures

5.13 (a) UV-Visible absorption spectra of sample SiC3, and (b) samples SiC7 and SiC9. 151

5.14 UV-Visible absorption spectra of samples SiC7 and SiC9 after calcination. 152

5.15 Chemical formula of epoxide group. 154

5.16 Chemical formula of DGEBA. 154

5.17 Reaction of Epoxy group with amine group. 155

5.18 The photograph of different dispersions just after ultrasonication at 65°C for 30 min

and after 24 hours of ultrasonication (1.Benzyl alcohol, 2.Benzene, 3.Isopropyl

alcohol, 4.Ethanol amine, 5. Toloune, 6. Chloroform, 7. Ethanol). 156

5.19 The photograph of pure DGEBA. 157

5.20 The photograph of nano SiC – DGEBA composites with increasing concentration of

filler from left to right (0.25%, 0.50%, 1%, 1.5%, 2% respectively). 157

5.21 SEM images of different composites, (a) 0% filler, (b) 0.25% filler, (c) 0.50% filler,

(d) 1% filler, (e) 1.5% filler, (f) 2% filler. 158

5.22 FTIR Spectra of pure DGEBA after treatment with NaOH, H2SO4 and NN-

dimethylformamide. 160

5.23 FTIR Spectra of nano-SiC- DGEBA composites with different filler concentration

before and after treatment with H2SO4. 161

5.24 TGA graphs of nano-SiC- DGEBA composites of different filler concentration. 161

List of Tables

List of Tables

1.1 Electrical properties of silicon [34]. 21

1.2 Thermal properties of silicon [34]. 22

1.3 Seven of the most simple SiC polytypes with four notations (R = Ramsdell

notation R, J = Jagodzinski notation R, Z = Zhadanov notation R). They are

listed by increasing percent hexagonality. 33

1.4 Hexagonalities, observed minimum indirect and direct bandgaps at 4 K and

their temperature dependences for various typical SiC polytypes [73-76]. 34

1.5 The optical modes and corresponding Raman frequencies of 3C and 6H

polytypes [76]. 35

1.6 Mechanical and electronic properties of 4H, 6H and 3C polytypes of SiC in

comparison with silicon and diamond [85], [86]. 36

3.1 The parameters of the DC arc Plasma reactor. 75

3.2 FTIR absorption peaks corresponding to different vibrations of bonds of Si

with O2, H2 and C [8,9]. 77

4.1 The details of the parameters used in different synthesis experiments. 92

4.2 The details of the synthesis parameters. 108

4.3 Comparative field emission study on silicon nanostructures. 124

5.1 The details of the synthesis parameters used for the synthesis of SiC-

nanoparticles. 133

5.2 The weight percent of impurities in SiC samples calculated from XRD pattern

and TGA. 140

5.3 The hardness values of SiC – Epoxy composites with increasing filler

percentage. 159

5.4 The percent weight losses of different composites. 162

1

Chapter 1 Introduction and Scientific

Background

This chapter gives an introduction to the research topics addressed in the thesis. Objective and

structure of the thesis is described. A brief scientific background about the method and the

materials used in this thesis is provided.

Chapter 1.Introduction and scientific background

2

1.1 Introduction

Nanotechnology is the potentiality to manouevre individual atoms and molecules to

produce nanostructured materials and sub-micron objects that have applications in the real

world. The exceptional properties of nanostructured materials have attracted researchers of

varied fields to embrace nanotechnology, which can be observed from steady growth in the

number of publications since its practical realization around 1990 (Figure 1.1 (a)). The data

was obtained from the web of science by searching the word ‘nano’. The numbers of

publications have increased almost exponentially with years. Figure 1.1 (b) shows the data

of the global investment in nanotechnology which again shows exponential increase [1].

Figure1.1 (a) Plot of number of publications per year obtained from ‘web of science’ on searching

‘nano’ (b) Public R & D investments in nanotechnology globally [1].

Due to the extensive research all over the world, it is now clear that nanomaterials

can be obtained mainly by two approaches; top down and bottom up. Top down approach

includes reducing the size of bigger to smaller e.g. ball milling, etching techniques, etc.

Bottom up approach includes growing nanostructures (NSs) from atoms and molecules e.g.

chemical synthesis, physical vapour deposition and plasma assisted synthesis techniques,

etc. Reducing size to nano regime mainly induces two major effects: surface to volume ratio

and quantum confinement effects. The first effect determines the physical, chemical and

electronic properties of the nanomaterials, which in turn vary considerably with the crystal

sizes. The second effect modifies the band structure of material. Thus, by precisely

controlling the size and surface of a nanocrystal, its properties like the bandgap, crystal

structure, electrical conductivity, melting point and different physical properties can be


3

tuned. The work presented in this thesis encompasses the bottom up approach of fabricating

nano crystalline structures of the well known semiconductor silicon.

The progress in semiconductor, especially silicon, has changed the world of

technology and has immensely contributed to the well being of mankind. The incessant

effort put into further miniaturization, governed by Moore’s law [2] demands investigation

of silicon nanostructures (SiNSs) in various forms. Apart from miniaturization, its optical

properties also require that research into Si be revisited. Even though Si dominates the field

of electronics, its photonic properties are considered poor due to its indirect band gap; this

band gap involves a momentum-balancing phonon (lattice vibration) during photon

absorption and emission. This disqualifies the use of bulk Si in optoelectronic devices when

compared to its competitor. The quantum confinement increases the probability of radiative

recombination [3,4]. So, different morphologies of SiNSs have been synthesized to achieve

the quantum confinement effect.

Nanocrystalline Si was synthesized way back in 1956 by Uhlir [5], but it was in 1990

when Canham [6] reported the visible PL from electrochemically etched porous Si. This

indirect evidence of free standing Si quantum wires and it's new light emitting properties not

only initiated but also gave a great momentum to the research in nanostructured Si. In spite

of the tremendous efforts, no commercial application using porous Si as the photonic

materials is viable due to the deteriorating optical properties, and the delicate structure that

cannot withstand the routine micro-fabrication technology. The other options like Si

nanocrystals embedded in silicon oxide, amorphous Si thin films, hydrogen-passivated Si

thin films, etc. were therefore looked into as a better option to porous Si and many of these

are being used for actual applications. Looking at the advantages and applications of Si it

has remained the favourite of researchers and there is consistency in publications related to

silicon research (Figure 1.2).


4

Figure 1.2 Plot of number of publications on silicon published per year (searched word silicon in

Science Direct).

Free standing NSs of Si are also important entities and have applications in varied

fields spanning medical, polymer, space, etc. Although different NSs of Si are widely

explored by now; silicon nanotubes (SiNTs) could not be studied extensively, due to

difficulty involved in its synthesis. The interest in SiNTs arises because the counter

element of silicon namely carbon has demonstrated the ease of the formation of two

dimensional sheet structure (graphene) as well as the folded sheets i.e. carbon

nanotubes (CNTs). The attractive properties of CNT and graphene have compelled the

scientific community to study these materials. Like CNT and graphene, “Why not

SiNT and silicene?” and “What would be their properties?” The topic of synthesizing

SiNTs, therefore, remains challenging.

After studying properties of Si, we feel that there should be a material which will

exhibit semiconducting properties of Si with higher breakdown voltage and can work till

higher temperature. Is there such material? The answer is yes and it is silicon carbide. SiC is

a compound of Si and carbon, which replaces Si for high temperature and high power

devices electrical devices. Due to a combination of desirable mechanical properties, such as

high hardness, wear resistance, strength at elevated temperatures in addition to corrosion

resistance, chemical inertness, electromagnetic response and bio-compatibility, advanced

ceramics are widely used for the manufacture of components for the optical, electronic,

mechanical and biological industries [7]. Furthermore, due to its low coefficient of thermal

expansion, high thermal conductivity, high decomposition temperature, low wettability by


5

molten metal and low density, SiC is commonly used for heat resistant parts and refractory

applications [8–10].

Of these, powdered SiC have been used in many sector, few to mention are ceramic

industry, grinding wheels, abrasive paper and cloths, as a high grade refractory material, in

manufacture of rubber tyres and heating element, in making of high temperatures sealing

valves, in modifying the strength of alloys and in mirror coatings for high ultraviolet

environments. Here, if the size of the particles is reduced to nano regime, the reduced

volume to surface area effect would become evident and the requirement of the material

would reduce by several percentages. In case of grinding and cutting tools the fineness of

the tools improvises. The density of ceramics and reaction bonded SiC will get enhanced.

Thus nanotechnology in SiC plays a very important role. However, SiC does not

occur naturally and require high temperature conditions for synthesis and obtaining

impurity free SiC is also difficult. Thermal plasma assisted technique is one of the high

temperature process which can be used for synthesis of SiC nanoparticles (SiCNPs).

In general, plasma processes [11,12], have proved to be effective tools for fabricating

different kinds of nanoparticles (NPs), thin films and coatings. Thermal plasmas, produced

by different routes, have been used for material processing, melting and production of high-

quality, high-crystalline materials as well as nanomaterials [13]. Due to their unique

advantages thermal plasmas are preferred in material processing; these include high enthalpy

to enhance reaction kinetics, high chemical reactivity, oxidation and reduction atmospheres

in accordance with the required chemical reactions, and rapid quenching to produce

chemically non equilibrium phases in nanomaterials. The high energy densities of plasmas

help in increasing the processing rates which lead to a large flux of radical species that are

crucial in the formation of metastable or transition phases [13–16] like SiNTs. Such

crystalline phases find applications in varied fields, due to the availability of large functional

sites. The added advantage of thermal plasma lies in its capacity to synthesize nanomaterials

on the large scale [14], required in industrial applications. These may include high energy

materials [15], catalysts [16], fabrication materials and those required in metallurgical works

like SiC. The interest in study of thermal plasma can be observed from the data (Figure 1.3).


6

Figure 1.3 Plot of publications and scitations (cumulative) with year by searching for ‘thermal

plasma’ (http://academic.research.microsoft.com/).

The work described in this thesis was carried out for understanding the importance of

SiNSs, SiCNSs and the role of thermal plasma to accomplish the synthesis of these NSs.

1.2 Objective of thesis

SiNTs are less commonly reported, the reason of which is accredited to the

preference of sp3

hybridization in Si which favors the formation of silicon nanowires

(SiNWs) compared to SiNTs. However, the theoretical models included tubular structures

built of hexagons of Si in both sp2

[17] or sp

3 hybridization [18]. Fagan et al. [17] have

estimated SiNTs to behave as semiconductors with a band gap of 2 3 eV while Bai et.al.

have predicted single-walled SiNTs to be metallic [19]. These properties of SiNTs are

significant and possibilities that they could be synthesized using thermal plasma are

predicted. Hence, the work is focused on the synthesis of SiNTs using arc plasma, study of

its structure and applications. It was aimed to study their field emission and investigate

biological applications. Antibacterial study for two strains of each gram positive and gram

negative bacteria is carried out.

Other material, on which the work is focused, is SiC. SiC is not a naturally occurring

mineral and it could be synthesized by processes involving high temperature conditions.

Thermal plasma assisted gas phase condensation is one of such routes using which SiC can

be synthesized. Most of the thermal plasma assisted processes involve use of gaseous

precursors like methane, silane, and H2 like ref [20]. Use of these precursors, add to

pollution and cost. Moreover, the product yielded consists of impurities of Si and carbon or

either of them [21]. The aim was to obtain SiCNSs by using solid Si and carbon precursors

without the use of gaseous precursors and H2. As an application area, the SiCNPs have been

http://academic.research.microsoft.com/


7

investigated in terms of its hardness properties by making composites with diglycidyl ether

bisphenol A (DGEBA) epoxy.

1.3 Organization of thesis

The thesis has been organized in six chapters as follows:

1. In this chapter, a brief background about the materials and the synthesis process used in

this work has been described.

2. In the second chapter, a thorough literature survey is presented about different SiNSs

their properties and applications as well as the theoretical predictions and methods of the

synthesis of SiNTs have been included. Similarly, literature survey about SiCNSs has

been presented.

3. Chapter 3 will deal with the experimental methods employed and a brief overview of the

characterization tools used for characterizing the as-synthesized NSs.

4. Chapter 4 includes the results of the synthesis of SiNSs. It consists of discussion about

the optimization of synthesis parameters to synthesize SiNTs, the detail analysis of

SiNTs and their field emission study and antibacterial activity.

5. Chapter 5 describes the results of synthesis of SiCNSs. It also consists of the discussion

about the fabrication and characterization of SiCNPs-DGEBA composites.

6. Finally, Chapter 6 will summarize and conclude the salient findings under the scope of

this work.

1.4 Scientific background

A) Synthesis method

1.4.1 Arc plasma

1.4.1.1 Definition

Arc is defined as a quasi-neutral ionized state of matter generated in between two

oppositely polarized electrodes sustained with the help of a constant input of electrical

energy. In general, the arc must be defined in terms of current and voltage drop only. It is a

class of electrical discharge where the current exceeds a threshold situated somewhere


8

between 0.1 and 1 A, the upper limit being unspecified and very large and the voltage drop

is in the range between a few volts and a few tens of volts.

1.4.1.2 Types of gaseous discharge

When a variable potential difference is applied between two electrodes separated by

a distance, different types of gaseous discharge are observed in different voltage and current

regime. These can be observed in figure 1.4 (a).

Initially, at low voltage, current density of the order of 10-14

A/cm2, are observed due

to the initial number of electrons present due to cosmic rays only. With increasing voltage,

current increases, then reaches a saturation current region and further increase in the energy

of electrons and their collision with gas atoms generate electron-ion pairs, which contribute

to further current. This onset of participation of both electrons and ions in the process of

charge multiplication through collisions is identified as townsend discharge regime which is

not self sustaining with typical current densities extending from 10-16

to 10-6

A/cm2.

Figure 1.4 (a) Current-voltage characteristic of a DC discharge through a gas [22] and (b) the

dependence of individual specie temperature in thermal plasma on pressure.

Further attempt, to increase current, results in a sudden drop in voltage with current

densities in the order of 10-4

A/cm2. This is the glow discharge regime, where discharge

becomes visible and gas is said to have undergone the process of electric breakdown.


9

Current carriers are now created and sustained by the action of positive ions and photons on

cathode and of electrons inside plasma hence it is a self sustained discharge. Major physical

processes contributing to glow discharge are mainly secondary emission and photoelectric

effects at the cathode. The characteristics of glow discharge can be divided into three main

regions namely; corona, normal and abnormal regions. Further increase in current density

i.e. from 10-2

A/cm2

to 10-1

A/cm2, requires an increase in voltage that is accompanied by an

increase in cathode glow tending to fully cover the cathode surface keeping current densities

constant. Beyond that is the beginning of abnormal glow discharge.

Further increase, in current, leads to sudden transition to a low voltage

discharge mode (~ tens of volts) known as arc discharge with current of the order of 1-

100 A or even higher. Arc mode is characterized by a violent discharge with intense

radiation emanating from the plasma and excessive heat loads to electrodes. At the

same time, arc current is restricted to the small areas on electrodes leading to very

high current densities (108-10

10 Am

-2) and thus, the heating of cathode surfaces [22]. As

a result electrons are emitted from the cathode due to thermionic emission process.

Calculations indicate that field emission cannot maintain an arc discharge independently and

is always accompanied to some extent by thermionic emission. Arc discharges can occur

over a wide range of pressures. The dependence of individual specie temperature on

pressure is illustrated in figure 1.4 (b). When the system pressure increases, the total

collisional frequency also increases and the temperatures of the individual species tend to

come closer. In the transition region one finds a slight drop in electron temperature and a

significant rise in the ion temperatures. Finally in the ‘high pressure arc’ (atmospheric and

above) all these species attain nearly the same temperature, thereby pushing a plasma into a

particular state of the plasma, called ‘local thermodynamic equilibrium’. This expression

means that, at each point of the plasma, one can define a unique temperature (and other

thermodynamic parameters depending upon the content). However, the equilibrium is purely

local; important heat fluxes, associated with significant temperature gradients, can exist

throughout the discharge.

1.4.1.3 Generation of arc

An arc is ignited when two electrodes, having a potential difference roughly a few

tens of volts, are brought very close to each other (of the order of few microns) in a gaseous


10

ambient. The surface of electrode consists of some sharp edges in the micron level. Thus,

the electric field at these sharp protrusions is very high that initiates field emission from

these edges and hence it ionizes the surrounding gas. When the two electrodes with applied

potential difference are approached to a small distance, the space charge produced at the

surface of the cathode causes sufficient field distortion to move the electrons ahead towards

anode. The rapidly moving electrons leave a tail of positive ions and generate new electrons

due to collisions. This process generates the avalanche of electrons that contribute to the

ionization of gas. The ionization of the ambient gas builds up electrically conducting

channels in between them where a considerable Joule heat is produced on account of the

transformation of the electrical energy into the heat energy. The amount of heat, thus

produced, decides the degree of ionization (the fraction of the species getting ionized) and

the overall enthalpy content of the arc and the gas is gradually turned into thermal plasma.

The method of arc generation and hence thermal plasma generation discussed here is by

electric discharge, which is used in the present study.

Other methods used for thermal plasma generation is by the use of electrode-less

discharges by radio frequency (RF), microwaves, shock waves, and laser or high-energy

particle beams. An RF discharge can be maintained either by capacitive or inductive

coupling with the power source. Capacitively coupled RF requires extremely high

frequencies for producing thermal plasma whereas inductively coupled RF discharge relies

on time-varying magnetic field and requires frequency in the range between 3 MHz and 30

MHz. Microwave discharges require frequencies ranging from 1 MHz to 10 GHz abide to

pressures from 10-3

Pa to several hundred kPa. Also, the plasmas can also be produced by

heating gases (vapors) in a high-temperature furnace [23].

After the initiation of arc, the ionization created needs to be supported by a suitable

power source which would supply power equal to the power dissipated in the plasma else it

would not sustain. Arcs can be sustained by direct current or alternating current with suitable

characteristics.

1.4.1.4 Regions of Electric-Arc

For low pressure as well as high pressure arcs, it is customary to distinguish three

regions; cathode region, anode region and arc column. The discharge appears constricted

both towards the anode and the cathode; although, in most of the cases, it is significantly


11

more constricted at the cathode end. This enables one to divide an arc into three regions, as

indicated in figure 1.5, which gives a schematic representation. The distribution of arc

voltage along the arc length is shown in figure 1.6

Figure 1.5 Schematic of the characteristic arc-regions in an unspecified electric-arc.

Figure 1.6 Schematic distribution of arc-voltage along an unspecified arc-length.

1.4.1.5 Building blocks of arc

(a) The Cathode

The flow of the current through the arc is affected by the electrons liberated by

thermal and field emission from the surface of the heated cathode. By acceleration in the

region of the cathode fall, the electrons acquire a high measure of kinetic energy, which

enables them to ionize neutral atoms in collisions. The positive ions, thus formed, are

accelerated in the opposite direction, strike the cathode and transfer their energy to it. Thus,

they heat the surface of cathode and keep up the thermal emission. No unified theory

explaining the cathode phenomena has been elaborated so far. Various theories are reviewed

in Ref. [24]. The cathode phenomena can be classified in three groups: arcs with a

constricted cathode region, arcs with an un-constricted cathode region and arcs with non-


12

stationary constricted cathode region. The current density of the cathode spot ranges from

108-10

10 Am

-2.

(b) The Anode

The anode is intensively heated by the electrons accelerated in the electric field

existing in the arc zone. During operation, the anode temperature may rise to the boiling or

sublimation point of the anode material. The atoms evaporated from the anode enter the arc

plasma and are ionized there. Unlike the cathode, however, the anode may be cold. The arc

burns well even with a water cooled anode and no ions of the anode material have been

found in the arc plasma in this case. During operation, a hissing of the arc can occur as a

direct consequence of the overloading of anode by the excess amount of current. As the

current in the arc is increased, the current density (in the micro-spots) of the anode rises to a

typical value of 108Am

-2 and a marked constriction appears in the anode region. The thermal

load in the micro-spots rises to extreme values (106

kcal m-2

). The eruptive escape of the

evaporated anode material is accompanied by a characteristic hissing sound and the anode

spot moves at a high velocity on the front face of the anode during the process. The

frequency of the hissing sound depends on the thermal conductivity of the anode and the

sound effect is closely connected with oscillations observed in the arc voltage.

On an average, the anode region can further be divided into four partial zones. The

arc column (1st) is followed by the transition zone (2

nd), where the directed motion of the

electrons becomes uniformly accelerated towards the anode. In acceleration zone (3rd

),

electrons acquire kinetic energy required for the ionization of neutral atoms, followed by

ionization zone (4th

) where the neutral atoms are ionized. The ion density is at its maximum

at the transition from the 3rd

to the 4th

zone and then decreases towards the anode, whereas,

electron density increases in this direction. The validity of theory is limited to low-current

densities and arcs with a high anode fall (~20V). At higher currents, electric field ionization

gives way to thermal ionization and the anode voltage may drop far below the value of the

ionization potential.

(c) The Arc Column

In the column of the arc, the gradient of the electric field is relatively low and its

magnitude is affected (apart from the current flowing through the arc) by various external


13

factors viz the kind of the gas, its pressure, the electrode material, the cooling of the arc and

external mechanical and magnetic forces.

As in the cathode and anode regions, the increase of the current flowing through a

freely burning arc reduces the electric field gradient in the column; the density of the

charged particles in the plasma is dependent on its temperature, which in turn, is determined

by the local energy balance of the arc. The energy supply by Joule’s heating has to equal the

energy losses due to thermal conductivity, radiation and convection. As the current

increases, the temperature of the plasma rises; hence its electrical conductivity increases and

the electric field gradient, therefore, diminishes [25].

The electric field gradient of a freely burning arc is a function of the electrical

conductivity of the plasma and consequently of its temperature. The heat losses are mainly

due to the conduction and the quantity of the energy, thus removed, is proportional to the

radius of the plasma column. As the current increases, so does the radius. At a certain

magnitude of the current two extreme states can occur: the cross-section of the arc is either

too large or too small. At excessive cross-sections, the temperature and hence the electrical

conductivity of the plasma is relatively low and the current transfer, therefore, requires a

high gradient. Small cross-sections produce high plasma temperatures and consequently

high electrical conductivity. Yet, because of the small cross-section, the current transfer

again requires a high gradient. The arc with the minimum gradient must be somewhere

between these two extremities. As per Engel and Steinbeck [26] this fact is in consequence

of the principle of minimum voltage. In the stable arc 0dT

dEor 0

dr

dE. The kind of the gas

greatly influences the arc voltage. Inert gases, such as Ar or He form no molecules and

require far less energy for their thermal ionization than do polyatomic gases, such as H2, N2,

O2 etc, which requires to be previously dissociated. The thermal properties of the gas exert a

greater influence on the arc plasma than does the ionization potential. The highest electric

gradient in the arc column occurs, where the plasma is formed by H2 (ionization potential

being 13.59 eV), although the ionization potentials of He, Ar and N2 are higher. Even gas

pressure in the medium surrounding the arc influences the arc voltage. The influence of the

electrode material is proportional to the amount of the material evaporating into plasma.

However, even a small material evaporated from the electrode may produce a marked effect


14

on the degree of ionization and the electrical conductivity of the plasma, provided, its

ionization potential is low.

1.4.2 Mechanism of thermal plasma assisted synthesis of nanoparticles

The most important issue, in the thermal plasma synthesis of nanomaterials, is

related to the nucleation and growth of the species from the vapor solid equilibrium. On

account of its high enthalpy, thermal plasma provides the necessary thermal energy to the

solid feed material and vaporizes it into its constituents atoms or molecules. Subsequently

the vapor diffuses out towards the periphery or the fringe of the plasma plume. Here it

encounters a steep temperature gradient and therefore the metal precursor spontaneously

undergoes vapor to solid phase transformation. Thus, supercooling to a temperature below

the melting point results into supersaturation, this is the driving force for nucleation. It is an

example of homogeneous nucleation. If growth is assisted by a catalyst or occurs or is

supported on a solid substrate it is an example of heterogeneous nucleation. The precursor

atoms and molecules consistently diffuse towards the nuclei and the growth occurs through

condensation of these molecules over the nuclei. Further, these NPs may collide with each

other resulting into the coagulation leading either to the coalescence into larger particles or

agglomeration into chains of NPs [27].

In view of the importance of plasma parameters in controlling the size and shape of

the NPs, during gas phase condensation, a brief review of the well documented [28]

thermodynamic features of nucleation and growth is presented in this section.

1.4.2.1 Homogeneous nucleation

Nucleation begins with the formation of solid embryo (or cluster) consisting of few

atoms. For a given volume of gaseous system consisting of precursor atoms, G1 and G2 may

be defined as the Gibb’s free energies before and after the formation of the embryo of radius

‘r’ (Figure 1.7 (a)). The change in the Gibbs free energy ΔGr (=G1 G2), upon the

spontaneous formation of spherical cluster of radius ‘r’ is given by

(1.1)

where, is the change in the volume free energy due to the formation of the solid

embryo (new volume) from the vapour phase. is the interface energy per unit area of solid-

vapour interface at the surface of the embryo. Below the melting point of the embryo


15

(because of supercooling) is positive, so that, the free energy change (associated with

the formation of a small volume of solid) has negative contribution due to lower free energy

of the bulk solid while there is a positive contribution due to the creation of solid/vapor

interface. It costs free energy to add molecules to the embryo until the radius reaches a

critical value . Those embryos, which have radius , will be unstable and will

dissolve into parent phase. Only those nuclei, which have radius above, can become

stable. The critical value is characterized by the maximum point in curve as

shown in figure 1.7 (b).

Figure 1.7 (a) Formation of the embryo: G1 and G2 are the Gibbs Free energies before and after the

formation of embryo, (b) the Free energy change associated with homogeneous nucleation with

radius r at different temperatures, (c) addition of atoms from the parent phase into the interface of a

critical nucleus and (d) the temperature dependence of nucleation rate I and growth rate U [28]

Eventually, both the critical radius and the critical free energy difference are

inverse functions of the extent of supercooling given by,

, (1.2)

, (1.3)

where, is the melting point and is the enthalpy of solidification

1.4.2.2 Rate of nucleation

If the system contains atoms per unit volume, the number of clusters that have

reached critical size is equal to


16

(1.4)

The addition of one more atom to each of these clusters will convert them into stable

nuclei which is also called the supercritical particle. This happens with a frequency and

the rate of nucleation is thus equal to

, (1.5)

where, is a complex function which depends upon the vibrational frequency of the atoms

around the nucleus, the activation energy for diffusion of atoms or molecules in the vapour

and the interface area of the critical nucleus. is the enthalpy of activation for diffusion

of atoms across the interface.

From equations (1.2) and (1.3) it is clear that the role of supercooling, in controlling

the process of nucleation is significant. Consequently, the greater the super cooling and the

smaller the critical radius less energy is needed to form it. Thus, in the presence of steep

temperature gradient at the edge of the plasma there is a high probability of homogeneous

nucleation.

The melting point of the precursor species also controls the process of nucleation.

Materials with high melting point may require larger critical radius that can sustain the

growth and thus pose greater difficulty in nucleation. On the contrary, substances having

low melting point will have ease in the nucleation process.

Now, if we consider the nucleation rate given by equation (1.5) it is seen to depend

on a number of parameters. However, significant contribution arises from . Large extent

of supercooling arising from the steep temperature gradients in the thermal plasma reduces

and thus the nucleation rate increases significantly. Also, the number density of

available reacting species is large due to the high enthalpy deposition of the plasma

into the feed material. On account of the inverse dependence of on Tm, metal having low

melting point will provide high nucleation rate and vice versa.

1.4.2.3 Growth

Growth is the increase in the size of the product particle after it has nucleated.

Growth usually occurs by the thermally activated jump of atoms from the parent phase to the

product phase. The growth rate ‘U’ can be expressed by,


17

, (1.6)

where, ‘r’ is the radius of the particle.

Both nucleation rate (I) and growth rate (U) pass through a maximum at some

intermediate degree of super cooling, as is indicated in figure 1.7 (d), since these are

thermally activated processes.

In the gas phase synthesis, initiated by thermal plasma, there is a constant source of

precursor atoms/molecules; however the maximum nucleation and growth rate are met at

different temperature zones where I and U would be optimized.

1.4.2.4 Crystalline phase and shape

The crystal structure of a nucleus depends on the temperature of formation as

described by its phase diagram. Further, epitaxial growth occurs which has similar

crystalline structure as that of the nucleus. Thus, the crystalline phase of the product

synthesized in the thermal plasma process by homogeneous nucleation, depends upon the

temperature regime where the nucleation occurs.

The nanocrystal, grown by homogeneous nucleation process, exhibits prominent

facets that depend on the surface free energy of the crystal-planes. The surface free energy ,

for a crystal is not isotropic, hence the growth rate depends on the crystalline directions.

According to Wulff construction [28], the crystal bounded by several planes A1, A2, etc with

energies , etc, will adopt a shape that satisfies the condition given by

= minimum (1.7)

The surface free energies of major planes in Si are related by γ100 < γ110 < γ111.

According to Gibbs-Curie-Wulf principle, a crystal growing under equilibrium conditions is

formed by faces with minimum γ hkl value due to gradual displacement of faces with

maximum γhkl value growing at maximum rates. Nanowires (NWs) are seen to have the

wire axis along (111), (112) or (110) directions depending on the growth conditions. On the

other hands the tubes are supposed to be formed by rolling of the planar sheets which can be

parallel to (111) planes but are stable when buckled.

Moreover in case of Si, different shapes of NSs are mainly governed by the point

group symmetry of the Si molecule in the crystalline structure. NPs and NWs of Si originate

from sp3

hybridized molecular structure whereas nanotubes (NTs) have preferentially


18

supposed to consist of a combination of sp3 hybridized buckled tetrahedrons and sp

2

hybridized planar Si structures, as discussed in section1.5.1.3. Such kinds of bond

configurations can be more probable because of steep temperature gradients (~ 104 K/cm)

and fast quenching resulting into the high degree of supersaturation in thermal plasma

systems. The crystal structure cannot relax in such a rapid process since the growth rate

would be faster than the relaxation rate and the nanocrystal gets frozen into a metastable

state. The shape of NSs is thus decided by the temperature gradients, rate of diffusion, and

ambient pressure during the nucleation and growth in the gas phase condensation.

B) Materials

1.4.3 Silicon

Si has become the most important and characteristic material of our age ‘the Si age’.

It has achieved this distinction with a rather modest volume of production as compared to

that of other basic industrial materials. There have been many attempts to find improved

materials with ‘better’ properties than Si, but candidates such as sapphire, SiC, diamond, II-

VI and III-V materials lack in: ease of growing large perfect crystals, freedom from

extended and point defects, existence of a native oxide, or other essential properties [29]. In

this sub-section, a brief discussion about properties of Si and SiNSs has been discussed.

Microstructure and properties of hypothetical SiNTs is also discussed.

1.4.3.1 Properties of silicon

Silicon is the second most abundant element found on earth (about 28% by mass)

and eighth most common element in the universe. It very rarely occurs in free elemental

form in nature and is mostly found in the form of silica and silicates. Si is a P Block, 14

Group, 3rd

Period element with atomic number 14, atomic weight 28.0855 and density 2.57

g/mL. Its electron configuration is [Ne] 3s2 3p

2. The 3s and 3p orbital undergo sp

3-

hybridization to form four equivalent orbitals with orientation in tetrahedral direction as

shown in figure 1.8.


19

Figure 1.8 sp3 hybridization in silicon.

These sp3 hybridized orbitals overlap linearly forming sigma bonds with bond energy

and bond length of 222 kJ/mol and 223 pm respectively. So, crystalline Si consists of

tetrahedra which are arranged in FCC diamond structure as shown in figure 1.9 (a). The sp3-

hybrid orbitals split into bonding and antibonding orbitals which constitute the valence and

conduction band.

Figure 1.9 (a) FCC diamond crystal structure and (b) the first Brillouin zone of the FCC lattice with

points of symmetry shown.

I. Band structure of silicon

The band structure of Si is shown in figure 1.10. There are three valence bands with

a single extremum at the center of the zone; the heavy hole band (Vl) with a hole mass of

mhh* = 0.46 m0, the light hole band (V2) with a hole mass of ml,h* = 0.16 m0 and a split-off

band with hole mass of mh,so* = 0.29 m0. The heavy hole band dominates the density-of-

states hole effective mass [30]. The minima at the bottom of the lowest conduction band

(C2) occur along the six principal cubic axes (along X, figure 1.9 (b)). The effective mass of

these anisotropic minima is characterized by a longitudinal mass (me,l* = 0.98 m0) along the

equivalent (100) direction and two transverse masses (me,t* = 0.19 m0, where m0 = 9.11 x 10-

31 kg) in the plane perpendicular to the longitudinal direction. The lowest band minimum at

k = 0 above the valence band edge occurs at Ec,direct = 3.4 eV. Thus, Si is an indirect-gap

semiconductor [31].


20

Figure 1.10 Schematic of the energy bands of silicon in the energy range near the forbidden energy

gap at 300K [31].

II. Optical transitions in silicon

Figure 1.11 (a) shows a schematic representation of direct and indirect optical

transitions in Si; strong optical transitions will occur for Δk = 0, at a higher energy than the

thermodynamic band gap at Δk 0.

Figure 1.11 (a) Schematic of direct indirect optical transitions in Si [32], and (b) absorption spectra

of single-crystal silicon at 77'K and 300'K [33].

Optical transitions occur at energies close to the thermodynamic band gap if both a

phonon and a photon are involved; emission or absorption of a phonon with the appropriate

wave-vector allows momentum to be conserved so that the transitions at Δk 0 are

possible. Figure 1.11 (b) shows the optical absorption coefficient of Si at 300 K and 77 K.

Both the indirect and direct transitions can be seen as absorption edges. The higher energy

direct transitions are stronger than the indirect transition. In addition, the indirect transition


21

is stronger at 300 K than at 77 K, as there is a substantial population of phonons of a suitable

wave-vector to take part in the transition at the higher temperature [32].

III. Electrical conduction in silicon

The electrical conduction in case of Si takes place by motion of electrons and holes

in a crystal, which is affected by collisions which change their speed or direction. These

collisions occur where the lattice periodicity is disturbed. Thus, a more perfect lattice at very

low temperatures would result in fewer collisions and greater mobility. Factors, such as

increased thermal agitation of the Si atoms, replacement of Si atoms in the lattice by

impurities, and the existence of crystal defects cause disturbances which can "scatter" the

carriers. The two most important scattering mechanisms in Si are lattice scattering and

ionized impurity scattering. The mobility of carriers limited by lattice scattering is

approximately proportional to T-3/2

, thus, this mechanism is dominant at higher

temperatures; while the mobility of carriers limited by ionized impurity scattering is

approximately proportional to T3/2

, thus, this mechanism dominates at lower temperatures.

Values pertaining to the electrical properties are mentioned in table 1.1, while other physical

properties of Si are mentioned in table 1.2.

Table 1.1 Electrical properties of silicon [34].

Property Value Units

Breakdown field ≈ 3·105 V/cm

Index of refraction 3.42 -

Mobility electrons ≈ 1400 cm2 / (V x s)

Mobility holes ≈ 450 cm2 / (V x s)

Diffusion coefficient electrons ≈ 36 cm2/s

Diffusion coefficient holes ≈ 12 cm2/s

Electron thermal velocity 2.3·105 m/s

Electronegativity 1.8 Pauling`s

Hole thermal velocity 1.65·105 m/s

Optical phonon energy 0.063 eV

Density of surface atoms (100) 6.78

(110) 9.59

(111) 7.83

1014

/cm2

1014

/cm2

1014

/cm2

Work function (intrinsic) 4.15 eV


22

Table 1.2 Thermal properties of silicon [34].

Property Value Units

Melting point 1687 K

Boiling point 2628 K

Specific heat 0.7 J / (g x °C)

Thermal conductivity [300K] 148 W / (m x K)

Thermal diffusivity 0.8 cm2/s

Thermal expansion, linear 2.6·10-6

°C -1

Debye temperature 640 K

Temperature dependence of band gap -2.3e-4

eV/K

Heat of: fusion / vaporization / atomization 9.6 / 383.3 / 452 kJ / mol

Energy of ionization:

Ist/IInd/ IIIrd/IVth

786.3/1576.5/3228.

3/4354.4

kJ.mol -1

1.4.3.2 Properties of silicon nanostructures

I. Quantum confinement effect

The most evitable effect of quantum confinement in Si is the widening of its band

gap when its size is reduced to few nm. Because of quantum confinement, valence states

shift down and conduction states shift up in energy; so the effective bandgap gets widened.

In a simple effective mass approximation the band gap shift due to quantum confinement is

given by,

,

(1.8)

where, m is the electron effective mass in the confinement direction, D is the diameter of

the potential well. Si nanocrystals with sizes in the range of approximately 5-40 nm show

size-dependent visible absorption in the range of 575-722 nm, while nanocrystals of average

size < 10 nm exhibit strong PL emission at 580-585 nm [35].

For NWs, the electronic properties, including band gaps, band structures, and

effective masses, are found to depend sensitively on all the nanowire structural parameters.

The size dependence of the band gap depends on the growth direction of the NW, and the

band gaps for a given size also depend on surface structure. NWs with reconstructed

surfaces have lowest unoccupied molecular orbitals (LUMO) localized on the reconstructed


23

facets and exhibit smaller band gaps. Effective masses of NWs, grown in the (001) direction,

decrease monotonically with size approaching the bulk transverse effective mass for large

wires. NWs with (011) growth directions exhibit a much weaker size dependence of the

effective mass [36].

Bulk Si has a very low quantum efficiency due to smaller lifetime for non radiative

transition τnr (~μs) than the radiative lifetime τr (~ns). Thus, it is unattractive as a light

emitting device [37]. The light emission in indirect bandgap Si nanocrystals can be

explained in terms of phonon assisted exciton recombination across the bandgap. For

radiative recombination, phonon must have the right momentum to bridge the separation in

momentum space between the top of the valence band and the bottom of the conduction

band. In bulk Si, thermal phonons ( kT ~ 26 meV) have enough energy to break-up the

exciton (energy about 15 meV) to a free electron and hole which move away from each

other through the continuum of states in the conduction and valence bands. Therefore,

radiative recombination becomes very unlikely as exciton break-up dominates. Whereas in a

nanoparticle, the continuum of the valence band and conduction band states is modified into

a discrete set of energy levels and the exciton binding energy increases due to the

confinement induced overlap of the electron and hole wavefunctions. In a Si quantum dot of

about 3 nm in diameter the exciton binding energy has been calculated to be larger than 160

meV [38]. Therefore, in a nanoparticle excitons cannot be broken up by thermal phonons,

thus, allowing the exciton enough time to wait for the phonon with the right momentum to

participate in the phonon assisted radiative recombination, producing an efficient light

emission at room temperature. Thus, light emission is observed in porous Si [39], Si

nanocrystals embedded in SiO2 matrix [40], silicon NPs [41] and NWs [42,43].

Reduction in the dielectric constant, refractive index and reflectivity arises owing to

the size effects. A similar lowering of the optical constants has been found for Si quantum

wells embedded in SiO2 [44]. All the calculations for Si nanocrystals yield a reduction of the

static dielectric constant which depends on the nanocrystal size. The reduction is significant

for nanocrystal size smaller than 2 nm [45].


24

II. Mechanical strength

Molecular dynamics with parallel computing technique was used to investigate the

effect of size on the mechanical properties of cubic SiNPs (side ~ 2.7 to 16.3 nm) [46]. The

results have indicated that the mechanical properties of the NPs are dominated by their

surface structures and the greatest maximum strength (24 GPa) is exhibited with a side

length of 10.86 nm. At lower values of the side length, the maximum strength reduces

significantly as the particle volume decreases. In the majority of cases, the maximum

strength and Young’s moduli of the current cubic SiNPs are significantly higher than the

equivalent values in the bulk system [46].

The surface Cauchy–Born model used to study SiNWs (diameters between 12 and 30

nm and aspect ratios between 8 and 32) shows that significant elastic softening is observed

in these NWs. The observed elastic softening does not manifest itself strongly until the

nanowire aspect ratio exceeds about 15. It is predicted by existing analytic models that the

elastic properties depend strongly on aspect ratio [47].

The experimental data of SiNWs (diameters ~ 40–90 nm, length ~several μm, outer

native oxide layer ~ 5 nm) shows that the nanowires can bear a large strain of 1.5% more

than the bulk material. The elastic constant of the nanowire was determined to be 175–200

GPa [48].

III. Surface texturing effect

Wang et al. [49] used the Bruggeman effective medium approximation along with

anisotropic optics to investigate the optical properties of SiNW arrays on Si substrates for

different polarizations at frequencies from 1 eV to 4 eV, which are of great importance for

solar photovoltaic applications. At low frequencies when the SiNW layer is semi-

transparent, the enhancement in the overall absorbance is mainly due to the antireflection

effect. At high frequencies, the SiNW layer with a few micrometer thicknesses becomes

opaque and can absorb more radiation than bulk Si. This is due to the dilution effect that

results in a smaller refractive index and subsequently a lower reflectance. The calculations

based on anisotropic optics clearly demonstrate the quasi-diffuse behavior of the optical

absorption for both polarizations; this results in improved hemispherical absorption over the

bulk Si. In the ideal case, the ultimate efficiency of the SiNW on Si substrate structure can


25

be as high as 42% at normal incidence and it is about 30% (relative) higher than that of bulk

Si. Yue et.al obtained the anti-reflection surface by fabricating hybrid structure of NWs and

pyramids, on the surface of silicon wafer by etching technique. The absorption of NWs was

found better over conventional pyramid textured Si by more than 10% [50]. Similarly,

SiNPs are coated on silicon solar cell to enhance its efficiency.

IV. Electrical properties

The electrical properties of the SiNPs depend on their size and shape, and

furthermore on the chemical composition of the protecting molecules. Some recent studies

describe the electrical properties of single SiNPs [51,52] as well as films of SiNPs

terminated with hydrogen or silicon oxide. The hydrogen passivated SiNPs show activation

energy of 0.49 eV. This is the energy necessary for an electron to hop from one nanoparticle

to the next neighbor in the film. For oxygen passivated NPs the activation energy increases

to 0.59 eV due to the isolating oxide shell. Functionalization with organic ligands generates

significantly higher activation energy in all cases like for n-octene and n-dodecene it is 0.66

eV and 0.68 eV, respectively, and is even significantly higher than that caused by the oxide

shell. For allylamine and allylmercaptan activation energy is 0.52 eV and 0.56 eV,

respectively. In a simplistic model, the activation energy for a hopping process from one

sphere to another is inverse simple proportional to the static dielectric constant of the

surrounding dielectric [53].

1.4.3.3 Silicon nanotubes

The interest in SiNTs has its source in CNTs. Although carbon and Si belong to

same group of periodic table, there is difference in their bonding character which results in

different morphology of their NSs. Theoretically, various atomic configurations of SiNTs

are assumed, and the structural stabilities and electronic properties are evaluated by diverse

calculational approaches. Here, the difficulty in the formation and theoretical predictions

about possible structure and properties of SiNTs, have been discussed.

I. Difficulties involved in the synthesis of silicon nanotubes

Carbon nanotubular structure shows efficient sp2 hybridization and p bonding

allowing formation of graphene sheet, thus high stability of the carbon nanotube structure


26

(Figure 1.12(a)). In contrast, Si prefers sp3 hybridization and favors the tetrahedral diamond-

like structures, thereby forming the commonly observed NWs (Figure 1.11(b)).

Figure 1.12 (a) Schematic of graphene sheet and carbon nanotubes and (b) schematic of stacking

observed in Si due to sp3 hybridization

The difference in the chemistry exhibited by carbon and Si can be traced to the

difference in their p bonding capabilities for which two components can be identified. First,

the energy difference between the valence s and p orbitals for carbon (is ΔE = E2p– E2s =

10.60 eV) is nearly twice that for Si (ΔE = E3p – E3s = 5.66 eV). As a result, Si tends to

utilize all three of its valence p orbitals, resulting in sp3 hybridization. In contrast, the

relatively large hybridization energy for carbon implies that carbon will activate one valence

p orbital at a time, as required by the bonding situation, giving rise, in turn, to sp, sp2and sp

3

hybridizations. Second, since the interatomic distance increases significantly in going from

carbon to Si, the p–p overlap decreases accordingly (by roughly an order of magnitude),

resulting in much weaker p bonding for Si in comparison with that for carbon. Hence, Si=Si

bonds are in general much weaker than C=C bonds [54].

II. Theoretically predicted silicon nanotubes and their properties

The structure of SiNTs is still an open question of fundamental physical and

chemical importance, which clearly requires rigorous efforts between theoreticians and

experimentalist. Different structures of SiNTs have been proposed by theoreticians. They

mainly include SiNTs formed of sp3-hybridized Si atoms, sp

2- hybridized CNT-like SiNTs

and sp2-sp

3 mixed type SiNTs.


27

a. sp2-hybridized SiNTs

The possibility of formation and stability of sp2 hybridized SiNTs have been

discussed by Fagan et al.[17]. The calculations also show that there is a significant cost

to produce graphite-like sheets of Si, but once they are formed, the extra cost to

produce the tubes is of the same order of the equivalent cost in carbon. The sp2

hybridized SiNTs with different chirality (zig-zag and armchair) studied by Barnard et al.

[55] are shown in figure 1.13. He has examined the importance of chirality and the diameter

on the structural, electronic and energetic properties of SiNTs. The calculations indicate that

the atomic heat of formation of a Si nanotube is dependent on the nanotube diameter, but

independent of the chiral structure of the tube. It has also been shown that the individual

cohesive and strain energies are dependent on both the diameter and chirality.

Figure 1.13 The figures of zig-zag (left) and armchair (right) silicon nanotubes structure [55].

The band-structure calculations by Fagan et al. [17] show that, similar to CNTs, the

band gap of SiNTs depends on the tube chirality. The tubes in armchair geometry show

metallic and the tubes with zigzag and mixed geometry show semiconducting behaviors.

Further, an ab-initio study of the energetic and structural properties of armchair and zigzag

SiNT structures, as a function of tube diameter is reported by Shan et al. [56]. The results

show that the band-gap properties are very sensitive to the deformation degree and the

helicity of the SiNTs.

Similar results were obtained by Durgun et al. [57] using ab initio molecular

dynamics calculations. Electronic analysis showed that zigzag NTs are metallic for very

small radii, but they show semiconducting behavior for larger radii while all armchair NTs


28

are metallic. They have also studied the mechanical behavior of SiNTs which show that the

SiNTs are radially soft; however they are strong against axial deformations.

The response of hypothetical SiNTs under axial compression is investigated using an

atomistic simulation based on the Tersoff potential. The results indicated that the application

of pressure, proportional to the deformation within Hook’s law, eventually led to a collapse

of the SiNT and an abrupt change in structure. Young’s modulus for SiNTs was constant

irrespective of the SiNTs’ diameter. As the SiNTs’ diameter increased, the collapse pressure,

that is the critical stress, linearly decreased [58]. Jeng et al. [59] found that the Young’s

moduli of the SiNTs are clearly lower than those of the conventional CNTs.

b. sp2-sp

3 hybridized SiNTs

Zhang et al. [54] have studied SiNTs using semiempirical molecular orbital PM3

method. Although, he describes difficulty in the formation of SiNTs, he has proposed that if

the dangling bonds are properly terminated, SiNT can in principle be formed. Tubular

structures for Si are, in general, less stable and tend to relax to the diamond-like structure

with tetrahedral configuration, which allows for the largest extent of overlap of the sp3

hybridized orbitals. Under appropriate conditions, partial structural relaxation and resulting

energy minimized SiNT, however, adopts a severely puckered structure (with a corrugated

surface) with Si Si distances ranging from 1.85 to 2.25 Å.

c. sp3-hybridized SiNTs

Metallic single-walled SiNTs have been reported by Jaeil Bai et al., based on the

calculations performed using molecular dynamic simulations. Model based on tetragons of

sp3-hybridized Si atoms for the possible existence of square, pentagonal, and hexagonal

single-walled SiNTs have been proposed (Figure 1.14 (a)). The local geometric structure of

these tubes differs from the local tetrahedral structure of cubic diamond Si, although the

coordination number of atoms is still fourfold. The calculations show that these tubes are

locally stable in vacuum and have zero band gap. Simillar structures of SiNTs (Figure 1.14

(b)) have been proposed by Lee et al. [60].


29

Figure 1.14 (a) The square, pentagonal, and hexagonal single-walled SiNTs proposed by Bai et al.

[19], and (b) the antiprismatic, prismatic and chiral SiNTs proposed by Lee et al. [60].

Bunder et al. [61] have proposed SiNTs formed by rolling up a two dimensional

quadrilateral lattice structure of sp3 Si atoms (Figure1.15 (a)). The quadrilateral lattice

shows some interesting band-gap behavior similar to hexagonal CNT lattices. They found

that the SiNTs are metallic in the majority of cases. Ponomarenko et al. [62] have studied

the energetics and relative stability of infinite and finite, clean and hydrogenated open-ended

SiNTs using the extended Brenner potential. The results suggest that the strain energy of

infinite SiNTs can be reduced by the chemisorption of atomic hydrogen onto the surfaces of

the tubes.

Figure 1.15 (a) An irregular quadrilateral lattice [61], and (b) side and front views of optimized

structures of infinite SiNTs clean (top), puckered with hydrogen capped on inner and outer surfaces

(middle), and ), with hydrogen capped on outer surface [62].


30

They have shown that the addition of hydrogen atoms to the Si hexagonal back-bone

of the tube to form SiNTs in which the Si–H bonds are oriented alternately on inner and

outer surface, reduces the strain energy and increases the chemical energy.

Seifert et al. [63] have considered silicides and SiH as precursors of possible sp3

hybridized SiNTs. The results show that such Si-based puckered NTs are indeed viable.

Using atomistic simulations within a nonorthogonal density-functional tight-binding

scheme, the structure, energetics, electronic and mechanical properties of silicide (111-Si

sheet) and SiH NTs have been obtained. The calculated values of Young's modulus of H-

terminated sp3 SiNTs are 7080 GPa. These values are over an order of magnitude lower than

for CNTs. Also these NTs are predicted to behave as semiconductors with a band gap of 2 -

3 eV. Varying the reaction conditions for silicide synthesis could therefore be a promising

way to fabricate silicide as well as SiH NTs. The authors also suggest the glow discharge

processes of monosilane as a possible way of SiNT synthesis.

Apart from these structures, Guo et al. [64] have proposed cluster-assembled

hydrogen passivated SiNTs (Figure 1.16). The results reveal that one-dimensional stable H-

SiNTs Sim(3k+1)H2m(k+1) can be built by stacking Si4mH4m cagelike clusters along the central

axis of the cage. Among all such SiNTs, the ones built from Si20H20 (m = 5) and Si24H24 (m

= 6) were found to be the most stable. The study indicates that hydrogen passivation may

be a good way to stabilize the hollow single wall SiNTs.

Figure 1.16 Cluster-assembled hydrogen passivated SiNTs proposed by Guo et al. [64].

d. Bulk silicon like SiNTs

These types of SiNTs include NTs with large wall thickness made up of bulklike

crystalline Si. Yan et al.’s [65] first-principles calculations for crystalline SiNTs show that

nonuniformity in wall thickness can cause sizable variation in the band gap as well as


31

notable shift in the optical absorption spectrum. The electronic wave functions of the

valence band maximum and conduction band minimum are mainly due to atoms located in

the thicker side of the tube wall. Based on the effects of nonuniform wall thickness on wave

functions of the valence band maximum and conduction band minimum of the SiNTs, they

have proposed a new modulation doping method, i.e., the selective p/n-type doping in the

thinner side to improve the carrier mobility and transconductance of doped nonuniform

SiNTs.

Chen et al. [66] have constructed SiNT structure by introducing a small hole at the

center and have found that thermal conductivity decreases for SiNT structure. The numerical

results demonstrate that a very small hole (only 1% reduction in cross section area) can

induce a 35% reduction in room temperature thermal conductivity. The enhanced surface-

to-volume ratio in SiNTs reduces the percentage of delocalized modes, which is believed to

be responsible for the reduction of thermal conductivity. Their study suggests that SiNT is a

promising thermoelectric material with low thermal conductivity.

1.4.4 Silicon carbide

SiC bears exceptional advantage because of its semiconducting properties added

with high thermal and chemical stability, and good hardness properties. Due to these

properties it finds applications in varied fields along with electronics industry. This sub

section describes the properties of bulk and nanocrystalline SiC.

1.4.4.1 Polytypism in silicon carbide

The basic unit of SiC consists of a covalently bonded tetrahedron of Si (or C) atoms

with a C (or Si) at the centre. The bonding of silicon and carbon atoms is 88% covalent and

12% ionic with a distance between the Si and C atoms of 1.89 Å. The identical polar layers

of Si4C (or C4Si) are continuously stacked and the permutation of stacking sequences allows

an endless number of different one-dimensional orderings without variation in

stoichiometry. These are known as polytypes. Figure 1.17 shows few of these polytypes

(2H, 3C, 4H and 6H).


32

Figure 1.17 Primitive hexagonal unit cells of the most simpleSiC polytypes. Si atoms are represented

by open circles, C atoms by filled circles. Bilayers of the three possible positions in projection with

the c-axis are labeled by the letters A, B, and C. The Si-C bonds in the (11 0) plane indicating the

relative shifts of the bilayers are represented by heavy solid lines. The figure has been reproduced

from K¨ackell et al. (1994) [67].

The Si-C bilayers are stacked on top of each other while they are laterally shifted by

1/ of the Si–Si or C–C atomic distance in the layer either in the or in the opposite

direction. If all shifts occur in the same direction, then an identical position of the bilayer in

the projection along the hexagonal axis is reached after three stacking steps. The resulting

structure is of cubic symmetry and because of the three-step stacking period this polytype is

called 3C (C for cubic) [68]. Another name for this polytype, which is the only cubic one, is

the often used term β-SiC. The other extreme is obtained, when the bilayers are shifted

alternatingly in opposite directions such that, in projection with the hexagonal axis, every

other layer has the same position. The lattice is then of hexagonal type, and because of the

two-step period the polytype is called 2H. All other polytypes are built up by a characteristic

sequence of cubic and hexagonal Si-C bilayer, for which the 3C and 2H polytypes represent

the limiting cases. All polytypes except 3C are uniaxial crystals (optical axis = c-axis) and

belong either to the hexagonal or to the rhombohedric system. The most abundant polytypes

besides 3C and 2H are the hexagonal types 6H and 4H and the rhombohedric 15R. The ratio

of the numbers of hexagonal to cubic bilayers is called hexagonality and is a very useful

scaling parameter. Several properties of the polytypes change with this parameter. Table 1.3

lists some of the simple SiC polytypes along with the four widely used notations and the

percentage of hexagonality.


33

Table 1.3 Seven of the most simple SiC polytypes with four notations (R = Ramsdell notation R, J =

Jagodzinski notation R, Z = Zhadanov notation R). They are listed by increasing percent

hexagonality.

R ABC notation J Z % of

hexa-

gonality

Space

group

No. of

atoms per

unit cell

3C ABC (k) ( ) 0 Td2(F 3m) 2

8H ABCABACB (kkkk)2 (44) 25 Td2(P63mc) 16

21R ABCACBACABCBACBCABACB (hkkhkkk)3 (34)3 29 C3v5(R3m) 14

6H ABCACB (hkk)2 (33) 33 C6v4(P63mc) 12

15R ABACBCACBABCBAC (hkhkk)3 (32)3 40 C3v5(R3m) 10

4H ABCB (hk)2 (22) 50 C6v4(P63mc) 8

2H ABAB (h)2 (11) 100 C6v4(P63mc) 4

Most of the polytypes, except 2H, are metastable. However, 3C does transform to 6H

at temperatures above 2000°C and other polytypes can transform at temperatures as low as

400°C [69,70]. The β-SiC (3C-SiC) with a zinc blende crystal structure (similar to

diamond), is formed at temperatures below 1700°C [71]. α-SiC (Wurtzite) is the most

commonly encountered polymorph; it is the stable form at elevated temperature as high as

1700°C and has a hexagonal crystal structure (similar to Wurtzite). Among all the hexagonal

structures, 6H-SiC and 4H-SiC are the only SiC polytypes currently available in bulk wafer

form.

1.4.4.2 Properties of silicon carbide

I. Band structure in silicon carbide

SiC is the only IV-IV compound to form stable and long-range ordered structures

polytypes). Over 100 different such polytypes have been observed. These polytypes are

semiconductors with a varying band structure. The energy band structures of SiC in the zinc

blende (3C-) and the wurzite structures (2H-SiC) have been calculated theoretically by many

authors since the first report in 1956 by Kobayashi [72].

From the theoretical calculations as well as the optical measurements, it is known

that all the polytypes have a valence band maximum at the zone centre (F-point). However,

the location of the conduction band minimum in k-space depends on the polytype, for

example, X-point for 3C- and K-point for 2H-SiC. All the polytypes, studied thus far, have


34

indirect bandgaps, which increase monotonically with the hexagonality of the polytypes h,

from Eg = 2.417eV for 3C-SiC (h = 0) to Eg = 3.33eV for 2H-SiC (h = 1) (Table 1.4).

Table 1.4 Hexagonalities, minimum indirect and direct bandgaps at 4 K and temperature

dependences for typical SiC polytypes [73-76]

Polytype

(Ramsdell)

% of hexagonality Minimum Bandgap (eV) dEg.ind/dT

Indirect Direct

3C 0 2.39 5.3 -5.8 X 10-4

8H 25 2.728

21R 29 2.853

6H 33 3.02 -3.3 X 10-4

33R 36 3.013

15R 40 2.986

4H 50 3.263

2H 100 3.33 4.39

II. Optical absorption

The optical absorption in SiC can, in general, be characterized by intraband and

interband absorption components. The interband transitions in n-type polytypes other than

3C are responsible for the well-known colours of nitrogen-doped samples. The intraband

absorption is of the free-carrier type and results in sub-bandgap transitions that are found in

most forms of SiC. Biedermann [77] has measured the optical absorption bands at room

temperature along the E c and E || c directions for 4H, 6H, 8H and 15R n-type SiC. These

are the most anisotropically used types of SiC. These polytypes are uniaxial and are strongly

dichroic [78]. The surface is normally perpendicular to the c-axis. This gives rise to the

green colour in 6H, the yellow colour in 15R, and the green-yellow colour in 4H polytypes.

These bands responsible for the colour are attributed to optical transitions from the lowest

conduction band to other sites of increased density of states in the higher, empty bands [79],

thus producing the various colours of nitrogen-doped 6H, 15R and 4H. Cubic SiC changes

form a pale canary yellow to a greenish yellow when the material is doped. The yellow

arises from a weak absorption in the blue region [79,80]. The shift towards the green for


35

doped 3C-SiC is due to the free-carrier intraband absorption, which absorbs red

preferentially [79, 76].

III. Phonons in SiC

Cubic SiC crystallises in the zinc blende structure (space group Td2(F 3m)), which

has two atoms per unit cell, and thus three optical modes are allowed at the centre of the

Brillioun Zone. Since SiC is a polar crystal, the optical modes split into one non-degenerate

longitudinal optical phonon (LO) and two degenerate transverse optical phonons (TO). 2H-

SiC, the rarest polytype, has the wurtzite structure and belongs to the space group

C6v4(P63mc). This uniaxial crystal has four atoms per unit cell, and consequently has nine

long-wavelength optical modes. Group theory predicts the following Raman active lattice

phonons, near the centre of the BZ: an A1 branch with phonon polarisation in the uniaxial

direction, a doubly degenerate E1 branch with phonon polarisation in the plane perpendicular

to the uniaxial direction, and two doubly degenerate E2 branches. The A1 and E1 phonons are

also infrared active, while E2 is only Raman active. Like 2H and other hexagonal polytypes,

6H-SiC belongs to the space group C6v4(P63mc), but has 12 atoms per unit cell (Table 1.4),

leading to 36 phonon branches, 33 optical and 3 acoustic. Therefore the number of phonons

observed by Raman Scattering is greater for 6H than for 2H-SiC, resulting in an additional

complication in distinguishing between several normal modes with same symmetry. The

optical modes and corresponding Raman frequencies of 3C and 6H polytypes are mentioned

in table 1.5 [76].

Table 1.5 Optical modes and corresponding Raman frequencies of 3C and 6H polytypes [76].

Polytype Modes Ref

E1

(cm-1

)

E2

(cm-1

)

TO(2)

(cm-1

)

LO(2)

(cm-1

)

TO(1)

(cm-1

)

LO(1)

(cm-1

)

TO(1) – TO(2)

(cm-1

)

3C -

-

-

-

-

-

-

-

796

796.2

972

972.7

0 [16,17,

30,10]

6H 766

768

788

789

788 964

967

797

796

970 9 [16,17,

30]

IV. Electromagnetic properties of SiC

SiC is considered to be one of the important microwave absorbing materials due to

its good dielectric loss to microwave [81]. In microwave processing, SiC can absorb


36

electromagnetic energy and be heated easily. A loss factor of 1.71 for 2.45 GHz at room

temperature and 27.99 at 695°C was calculated by Zhang et al. [82]. This ability for

microwave absorption is due to the semiconductivity of this ceramic material [82].

Moreover, SiC can be used as microwave absorbing materials with lightweight, thin

thickness and broad absorbing frequency. Since pure SiC posses low dielectric properties

that gives barely the capacity to dissipate microwave by dielectric loss, therefore, doped SiC

was used in order to enhance the aimed properties. SiC shows absorption in the frequency

range of 8.2-12.4 GHz [83].

V. Other properties of SiC

SiC is a wide band-gap indirect semiconductor with high breakdown voltage and

high saturation electron drift velocity. It is chemically inert with high hardness. While all

polytypes of SiC exhibit quite similar mechanical and thermal properties, their electrical and

optical properties differ greatly from polytype to polytype [84]. The properties of different

types of SiC are mentioned in table 1.6 along with silicon and diamond for comparison.

Table 1.6 Mechanical and electronic properties of 4H, 6H and 3C polytypes of SiC in comparison

with silicon and diamond [85], [86]

Properties Silicon Polytypes of Silicon Carbide Diamond

4H-SiC 6H-SiC 3C-SiC

Melting Point (ᴼC) 1420 2830 2830 2830 4000

Density (g cm-3

) at 300 K 2.3 3.21 3.21 3.21 3.5

Thermal Conductivity (W

cm-1

K-1

) at 300 K

1.31 4.9 4.9 3.7 20

Thermal Expansion

Coefficient (K-1

)

2.6 10-6

- 4.5 10-6

3.0 10-6

0.8 10-6

Moh’s Hardness 7 9 9 9 10

Bulk Modulus

(dyn cm-2

)

0.97 1012

2.5 1012

2.2 1012

2.2 1012

4.4 1012

Band Gap (eV) 1.12 3.26 3.2 2.4 -

Dielectric Constant (k) 11.9 9.7 9.66 9.72 -

Breakdown field (MV/cm)

(at donar impurity 1017

cm-3

)

0.3 3.0

( to C-

axis)

3.0 ( to C-

axis)

>1 ( to C-

axis)

> 1.5 -


37

1.4.4.3 Properties of SiC nanostructures

Bulk SiC shows weak optical emission at room temperature [87] that can be

significantly enhanced when the crystallite size is reduced to several or tens of nanometers

[6,88]. This is thought to be caused by depressed non-radiative recombination in the

confined clusters [89]. In accordance with the quantum confinement (QC) effect,

photoluminescence (PL) of the crystallites with diameters below the Bohr radius of bulk

excitons is shifted to blue with decreasing sizes [90]. Theoretically, the structure and

electronic properties of SiC NSs have been investigated employing semi-empirical and first-

principle calculations [91–93]. The results suggest that the band gap of SiC is dependent not

only on the sizes and but also on surface compositions of NSs strongly. Quantum Monte

Carlo calculations show that the C-terminated and H-rich quantum dots have the largest

gap [94,95]. Wu et al. have observed a blue band in the 3C-SiC NSs owing to the quantum

confinement effect and an additional PL band at 510 nm when the excitation wavelengths

are longer than 350 nm. The 510 nm band appears only in acidic suspensions but not in

alkaline ones and were found to arise from structures induced by H+

and OH- dissociated

from water and attached to Si dimers on the modified (001) Si-terminated portion of the NCs

[96].

Fan et al. [97] have also studied the effect of solvents on photoluminiscence property

of SiC and found that the solvent in which the 3C-SiC nanocrystals are suspended plays two

critical roles i.e. it serves as the sustaining medium that keeps the individual 3C-SiC

particles apart and it provides a high potential barrier for the carriers (electrons and holes)

to ensure quantum confinement [97]. The theoretical calculations done using an infinite

square-well potential for the 6H-SiC crystallites show that particles smaller than 3 nm

exhibit significant band-gap widening, and this effect is minimal in particles between 4 and

7 nm. Moreover, the high chemical and thermal stabilities [98] of SiC make the

luminescence from these nanocrystals very stable enabling the use of the materials in harsh

environments [99].

Field emission properties of NSs of SiC especially NWs are well explored and they

are found to be good field emitters because of the small curvature of the tip radius, high

aspect ratio, chemical inertness, and electrical conductivity. SiC/SiOx nanocables synthesized

by thermal evaporation of carbon powders and silicon powders in the presence of Fe3O4


38

nanoparticle catalysts show the low turn-on and threshold electric fields of 3.2 and 5.3 V/m

at the vacuum gap of 200 m, respectively. When the vacuum gap was increased to 1000 m,

the turn-on and threshold electric fields were decreased to 1.1 and 2.3 V/m, respectively

[100]. Field emission examinations of oriented silicon carbide nanowires (SiCNWs),

synthesized using CNTs possess large field emission current densities at very low electric

fields (2.5–3.5 V/μm) [101]. The turn-on field of carbon-coated SiCNWs at the emission

current density of 10 mAcm-2

was about 4.2 V/μm [102]. The aligned SiCNWs are good

field emitter material.

With a decrease in the grain size, the hardness of SiC and other refractory

compounds substantially increases [103,104]. Thermal stability of SiCNPs is used to

increase the thermal stability of SiBCN ceramics. The permittivity, dielectric loss and

absorption coefficient of ceramics increased as an elevated SiC content, resulting from the

increase of carrier concentration. SiCNPs also increased the permittivity and dielectric loss,

indicating their great potential as the high-temperature microwave absorption materials

[105]. Besides, SiCNTs are predicted to bear better hydrogen storage capacity than CNTs

[106].

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43

Chapter 2

Literature Survey

This chapter consists of a brief review of literature regarding different synthesis techniques and

applications of silicon and silicon carbide nanostructures.

Chapter 2. Literature Survey

44

This chapter contains literature survey about synthesis techniques and applications of

the nanostructures (NSs) studied in this work. Silicon nanostructures (SiNSs) are very well

studied materials and reviewing all the literature about it would require writing books

together, here very briefly the important synthesis techniques of different SiNSs and their

applications in various fields have been summarized. Later, the reports of synthesis of

silicon nanotubes (SiNTs) and their applications are reviewed in detail. In the subsequent

section, the synthesis methods and applications of silicon carbide nanoparticles (SiCNPs)

are reviewed in brief followed by introducing the role of thermal plasma in its synthesis. The

literature pertaining to study of properties of NSs discussed in Chapter 1 has not been

explicitly described and included in this chapter.

2.1 Silicon

Bulk silicon being indirect band gap semiconductor bears disadvantage of low

quantum efficiency of emission. The observation of luminescence from porous silicon (PS)

in 1990 [1] was sought as a beginning of new silicon era. The synthesis of PS can be

accomplished by chemical route [2], electrochemical route [3], chemical dissolution

involving HNO3, NaNO2 or CrCO3 in HF [4], spark erosion of silicon substrate [5,6], etc.

PS bear disadvantages of the fragile mechanical structure due to the highly porous

nature and degradation of luminescence with time. Thus, silicon nanocrystallites in the

SiO2 matrix were investigated as an alternative to PS. Si rich SiO2 samples are grown by

chemical vapour deposition (CVD) [7], laser ablation [8], sputtering [9] or formation of

Si/SiO2 multilayers followed by annealing [10], metalloorganic CVD [11], ion implantation

[12] , or laser pyrolysis [13].

Other than the good optical emission the major deficit of bulk Si is in photovoltaic

devices. Due to phonon assisted transition the efficiency of crystalline silicon is limited to

the Shockley-Queisser limit [14] and growing of crystalline Si solar cells is a costly affair.

Thus, amorphous Si was being used for solar cells. Amorphous Si films can be fabricated

using plasma enhanced CVD [15], hot wire [16], photo CVD [17] and sputtering techniques

[18]. The main merit of amorphous Si is not efficiency, but the cost and that it can be grown

at lower temperatures. However, to increase its efficiency, the multi-layer construction is

required which again adds to the cost. So, the search for new technology is on.


45

Later, the answers to the drawbacks of Si begin to be expected from nanotechnology.

This led to the synthesis of the different forms of NSs, employing different methods. Strong

fluorescence was observed in silicon nanocrystals preferably with dimension less than 5 nm

[19,20]. Different SiNSs were investigated and with time they found applications in varied

fields. Thus, here few of these nanostructures and their applications have been reviewed.

2.1.1 Synthesis methods and applications of silicon nanoparticles (SiNPs)

SiNPs have been synthesized by various methods few of which are solution-phase

reduction [21,22], microemulsion [23,24], sonochemical synthesis [25], mechano-chemical

synthesis [26], laser ablation [27], plasma-assisted aerosol precipitation [28,29],

electrochemical etching [30–32], green synthesis [33,34] and microwave-assisted synthesis

[35,36], etc. Kauzlarich and coworker’s report of solution-phase reduction synthesis strategy

capable of mildly producing SiNPs at room temperature and normal atmospheric pressure

[21], Kortshagen et al.’s [37] report on plasma assisted aerosol precipitation for synthesis of

SiNPs with controllable sizes ranging from 2 to 8 nm and Lee et al.’s [32] report of

electrochemical etching method, allowing fabrication of multicolor fluorescent SiNPs with

tunable maximum emission wavelengths from 450 to 740 nm were few of novel works to

be mentioned.

Passivation of SiNPs

SiNPs were intended to apply as fillers in inks for the fabrication of printable

electronic devices. However, the particles must be long term protected against oxidation in

ambient air. So, different methods of passivation and modifying the surface of nanoparticles

have been reported. These include alkyl terminated SiNPs [38,39], siloxane coated [40],

amine terminated [41], hydrogen capped [42], styrene passivated [38] etc. The research on

chemical reactions of molecules attached to the surface of silicon quantum dots that have

been performed to produce quantum dots with reactive surface functionalities have been

reviewed by Shiohara et al. [24].

Bioapplications of SiNPs

Fluorescent SiNPs are highly promising for biological and biomedical applications,

due to favorable biocompatibility and low toxicity [43–45]. The biological and biomedical

applications include bioimaging, biosensor [46,47], drug delivery [48,49], etc. Details about

the biological applications have been described in the book by Yao He and Yuanyuan Su


46

[50]. However, SiNPs need to be dispersible in water for these applications. Thus, require

additional post-treatment or surface modification to render the prepared SiNPs hydrophilic

for biological and biomedical applications. He et al.51

have developed novel microwave-

assisted strategies to facilely and directly synthesize highly fluorescent and water-dispersed

SiNPs in aqueous phase [35,51]. Another method which gives water dispersable SiNPs is

green synthesis [33,34].

Besides larger SiNPs are also synthesized and studied for catalytic [53], silicon ink,

high refractive index composites [54], etc.

2.1.1.1 Thermal plasma assisted synthesis of silicon nanoparticles

Thermal plasma is vital for synthesizing nanostructures of covalently bonded

elemental semiconductors such as Si, Ge and compound semiconductors like III-V

compounds that require high temperature to produce the crystalline state. Moreover, the

photoluminescence (PL) efficiency of crystalline SiNSs greatly exceeds those of amorphous

and defect containing nanostructures. Thus, growing highly crystalline SiNSs is important.

Laser induced plasma produced with silicon vapor was examined by Cowpe et al.

[55]. The temporal evolution of the laser ablation plumes in air at atmospheric pressure as

well as at a pressure of 10-5

mbar is presented by the authors. Temperature measurements by

plasma emission spectroscopy showed that the electron temperatures range between 7600-

18200 K for the atmospheric plasma and 8020-18200 K for the low pressure plasma.

Electron densities in the range of 6.91×1017

to 1.29×1019

cm−3

at atmospheric pressure and

1.68×1017

to 3.02×1019

cm−3

under vacuum were observed. These measurements prove the

existence of thermal plasma conditions when laser induced evaporation is used for the gas

phase synthesis.

Dynamics and time evolution of SiNPs formation during laser ablation of silicon

target in argon atmosphere have been investigated by Makimura et al. [56]. They have

observed the growth of SiNPs on a time scale of 1 ms. It was also found that the NPs emerge

just after the thermalization of the ablation plume, growing just above the ablation spot and

slightly apart from the target. Both, the ambient gas pressure and the ablation-laser’s energy

density were found to be the important factors affecting the time scale of nanoparticle

formation and growth. In particular, nanoparticle growth is delayed at higher-energy density

or lower gas pressure.


47

The synthesis of silicon nanopowders has been investigated by Leparoux et al. using

an ICP process in which they used Commercial microscale Si powder as a precursor.

Nanopowders with specific surface areas varying from 69 to 194 m2g−

1, corresponding to

equivalent particle sizes of 37 and 13 nm respectively, could be produced [57]. The

precursor evaporation process in an ICP system for nanoparticle production has been

modelled adopting different models for turbulence and particle evaporation by Colombo et

al. [58]. Schreuders et al. have used computational fluid dynamics (CFD) to improve an

Inductively Coupled Plasma (ICP) process for nanoparticle synthesis. The influence of

several quenching parameters (e.g. flow rate, composition, and quench design) on the

particle size has been investigated by calculations and experiments [59].

Single crystalline nanocrystals of silicon were grown as a byproduct during the

electrical-discharge-machining (EDM) by Davila et al. [60]. The average size of NPs,

formed during the EDM process, was around 500 nm. This again is an example of gas phase

synthesis resulting from electrical discharge induced evaporation in the pressure of de

ionized water.

So et al. [61] have used silane in RF thermal plasma reactor to synthesize spherical

and well crystallized SiNPs. They found that the size of the NPs depends on input power,

quenching gas flow rate and carrier gas flow rate. The smallest mean particle size of 36 nm

was obtained from the highest carrier gas flow rate of 45 L/min.

2.1.2 Synthesis methods and applications of silicon nanowires (SiNWs)

First report of synthesis of SiNWs

SiNWs bear advantages of the excellent electronic/mechanical properties, huge

surface-to-volume ratios, facile surface modification, and compatibility with well developed

silicon technology along with the PL properties observed due to 2D quantum confinement

effect [62,63]. First preparation of Si whiskers with <111> orientation with macroscopic

dimensions was carried out by Treuting et al. in 1957 [64], followed by enlightening work

of Wagner and Ellis on the vapor-liquid-solid (VLS) mechanism of the Si whisker growth

[65] which opened exciting avenues for fabrication of SiNWs. In VLS method, certain metal

impurities is an essential pre- requisite for growth of SiNWs acting as a preferred sink for

the arriving Si atoms or, perhaps more likely, as a catalyst for the chemical process involved

[66].


48

SiNWs on substrate

Nanowire arrays grown on substrate are important for photovoltaic applications,

sensors, photonic devices, rechargeable batteries, etc [67]. Silicon nanowire textures on the

surface decrease reflection of the radiation depending on the size, which is important for

performance enhancement of many optical devices such as solar cells and planar displays

[68]. The intensity of work done in this field is evitable from the number of review articles

published in this field [67, 69–72].

Different methods of synthesis of SiNWs

Methods used for the synthesis of SiNW can be categorized as top down and bottom

up approaches. The bottom up approaches includes VLS - CVD [73–75], oxide-assisted

growth [76,63], molecular beam epitaxy [77], laser ablation [78], etc. Specifically, due to

elegant work of Lieber, Lee and Yang et al., CVD and oxide assisted growth [66,73,76]

have been widely employed as two most popular means to fabricate SiNWs and SiNW

arrays with high aspect ratio and production yield. Top-down technologies use of

nanofabrication tools such as e-beam lithography [79–80] lithographically patterned NW

electrodeposition, nano-stencil lithography [81], or nanoimprint lithography [82]. It also

includes etching techniques like metal-catalyzed electroless etching [83–85], plasma etching

etc. Peng et al. developed a class of metal-catalyzed electroless etching approach (e.g., HF-

etching-assisted nano-electrochemical method) [83–85], serving as an alternative method to

facilely produce SiNWs in a low-cost manner. Horizontal SiNWs are mostly fabricated from

either silicon-on-insulator wafers [86–87] or bulk silicon wafers [88] using a sequence of

lithography and etching steps, often employing electron-beam lithography and reactive ion

etching. There are excellent articles of Singh et al. [89] and Suk et al. [88] on these

techniques.

Bioapplications of silicon nanowires

SiNWs are widely used in biological applications. The surface of the SiNWs has to

be modified with a probe molecule so that the biosensor is capable of recognizing a specific

target molecule. In order to attach the probe molecule on the surface, two approaches,

electrostatic adsorption and covalent binding, have been mainly adopted for use of SiNWs in

biosensors. Park et al. [90] reported a novel approach to selectively functionalize the SiNWs

surface using joule heating, which is extremely valuable for nanowire-based sensor


49

developments [90]. Gold nanoparticles-decorated SiNWs are used as highly efficient near-

infrared hyperthermia agents for cancer cells destruction [91]. SiNWs are considered as

promising candidates for fabrication of high-performance biosensors [92, 93].

2.1.2.1 Thermal plasma assisted synthesis of silicon nanowires

Thermal plasma is mainly applied in synthesis of SiNWs via plasma assisted VLS

mechanism like Qin et al. have synthesized SiNWs using ICP - CVD. The nanowires

consisted of crystalline core surrounded by a thick amorphous silicon shell. Increase in

plasma power produced dense and long nanowires with thick amorphous shell, accompanied

with a thick uncatalyzed amorphous silicon film on the silicon substrate. An enhanced

optical absorption was observed due to the strong light trapping and anti-reflection effects in

the thin and tapered SiNWs with high density.

The thermal plasma assisted synthesis has been carried out by few researchers.

Lamontagne et al. have synthesized SiNWs from carbothermic reduction of silica fume in

RF thermal plasma. They have discussed the impact of the addition of catalysts and the use

of different plasma gases on the yield and the properties of the product using different

characterization techniques. They observed that metal catalyst promoted the formation of

SiNWs and improved the yield of the reaction upwards of 300%.

Direct current (DC) arc discharge method was used by Feng et al. [94] for the

synthesis of SiNWs. The SiNWs had homogeneous diameters of 10–20 nm and lengths

ranging from several ten nanometers to several microns. They observed that the morphology

control of the products can be easily achieved by adjusting the current and the voltage of the

discharge. The formed NWs were polycrystalline.

Korpinarov et al. [95] have reported the synthesis of SiNWs and nanowhiskers by

DC arc discharge using a graphite cathode and a graphite anode filled with Si and C powder

mixture. The reactor was operated in a pre-evacuated argon-filled chamber at a pressure of

3x104 Pa. The arc current was maintained at 75 A by a DC power supply.

Review articles and books on silicon nanostructures

In addition to SiNPs and SiNWs, silicon-based nanohybrids featuring

multifunctional properties are promising as powerful tools for various applications [44].

Besides, many researchers have reviewed the work in different field of SiNSs, which


50

includes review articles by Misra et al.[96], Xiu et al. [97], Koshida and Matsumoto et al.

[98], Adamo et al. [99], Kang et al.[ 100], Veinot et al. [101], Schierning et al. [71], Shao et

al. [67], etc. Also, there are several books on SiNSs entitled ‘Silicon Nanocrystals:

Fundamentals, Synthesis and Applications’ [102], ‘One Dimensional Nanostructures’ [103],

‘Silicon Nano-biotechnology’ [50], ‘Nanosilicon’ [104], etc.

2.1.3 Synthesis Methods and Applications of Silicon Nanotubes (SiNTs)

Theoretical investigations in SiNTs

Inspite of difficulties in the formation of SiNTs, a lot of theoretical work has been

carried out about the type of SiNTs and their expected characteristics. Different types of

SiNTs have been proposed and their properties have been studied like SiNTs formed of sp3-

hybridized Si atoms [105–108], sp2- hybridized CNT-like SiNTs [109–114] and sp

2-sp

3

mixed type SiNTs [115]. Even bulk silicon like SiNTs have been studied [116,117].

Single walled SiNTs proposed by Bai et al. [118] showed metallic nature. Serhan

Yamacli investigated the voltage dependent transport properties of metallic SiNTs [119].

Bogdan et al. [120] have investigated the vibrational properties of pentagonal and

hexagonal single walled SiNTs by using Density Functional Theory and the frozen phonons

method. Motohike Ezawa [121] studied the buckled (sp3 hybridized) SiNTs and found that

the buckling along the application of electric field leads to the tuning of band structure.

Andriotis et al. [122] found that the encapsulation of metals (Ni and V) could

stabilize SiNTs. They found that these metal encapsulated SiNTs were metallic in nature.

Density functional theory [123] based calculations show that metal encapsulation turns

SiNTs into a metal or a semiconductor with very small band gap. Chandel et al. [124] have

studied the differences in structures and bonding arrangement of encapsulated and

functionalized (6,6) SiNTs and their relative stabilities in terms of their characteristic

electronic structures on interaction with mono-atomically thin metal wires of Ag, Au and Cu

from within (encapsulated) and outside (functionalized).

Hang et al. [125] observed that in comparison to other SiNTs, armchair nanotubes

are more stable due to the sufficient overlap of pz orbitals and delocalization of π bonds.

Moreover, SiNT (6,6) was found to be more stable in comparison to (4,0) and (5,5) config-

uration [112,126]. Not only stability but also doping and functionalization of SiNTs are well

studied. Singh et al. [127] found that the band structure of the Mn-doped ferromagnetic


51

single walled hexagonal SiNTs (sp2-hybridized) showed a gap just above the Fermi energy

for the one spin component, which shows there could be possibilities of making half metallic

NTs by including a small shift in the Fermi energy. Mahmoud Mirzaeie has studied the

formation of boron and nitrogen doped sints using density functional theory calculations

[128]. Functionalization of NT's has also attracted a considerable interest in the fields of

physics, chemistry, material science and biology [129]. The functionalization of SWNTs

with biological molecules is a relatively new direction in exploring the chemistry of SWNTs

for biosensor applications [130]. Mirzaei and Meskinfam [131] have used density functional

theory calculations to investige nuclear magnetic resonance properties of (5,5) armchair and

(8,0) zigzag models of SiNT. Junga Ryou et al. [132] have studied hydrogen storage

property of hexagonal SiNTs. Zhang et al. [133] have found that all silicene-like nanotubes

are stabilized by endohedral metals.

Raad Chegel and Somayeh Behzad [134] have investigated the electronic properties

of SiNTs produced by rolling up a hexagonal sp2

and sp3 Si sheet, under the external electric

field, using tight binding approximation. Zhu et al. [135] have studied multiwalled SiNTs

using first-principles density functional theory calculations and molecular dynamics

simulations and showed that the interaction between the walls is preferable through

covalent bonds rather than weak Vander Waal interactions; CNT-like SiNTs do not show

good stability.

Experimental investigations in SiNTs

The most common nanotubes that have been synthesized consist of bulk like SiNTs.

These NTs have been synthesized by etching techniques, CVD, catalytic RF plasma

treatment, molecular beam epitaxy, galvanic displacement reaction, electrodeposition and

template methods or a combination of these techniques. Mbenkum et al. [136] have grown

SiNTs with a wall thickness of approximately 4-5 nm from quasi-hexagonally ordered gold

(Au) nanoparticle arrays on SiOx/Si substrates using CVD technique. Chen et al. [137]

prepared novel SiNTs with inner-diameter of 60-80 nm using hydrogen-added

dechlorination of SiCl4 followed by CVD on a NixMgyO catalyst.

Using alumina template

Most of the methods adopted for SiNT syntheses involved use of template method.

For example, ferromagnetic SiNTs were grown in a hot-wall CVD setup using commercial


52

anodized aluminium oxide (AAO) membranes as a template by Shpaisman et al. [138]. The

wall thickness of about 10, 17, 25, and 40 nm was obtained with Ni homogeneously

distributed over the whole nanotube. Sha et al. [139] have synthesized SiNTs by CVD

Process using a nanochannel Al2O3. Co nanoparticles assisted growth of SiNTs on the pore

walls of the AAO was carried out by Zhang et al. [140]. Jeong et al. [141] have synthesized

SiNTs on porous alumina using molecular beam epitaxy. Meng et al. [142] have synthesized

heterojunctions between NTs and NWs by a combinatorial process of electrodepositing

NWs within the branched nanochannels of AAO template, selectively etching part of the

deposited NWs, and growing NTs in the empty channels on the ends of the NWs.

Using nanowire as template

NWs of different types were also used as templates. Zhou et al. [143] used aligned

suspended polyvinyl pyrrolidone nanofibers array as template to obtain ultralong (~4 mm)

SiNTs by a hot wire CVD process. Mirko Battaglia et al. [144] synthesized amorphous

SiNTs grown in a single step into a polycarbonate membrane by a galvanic displacement

reaction conducted in aqueous solution. He also found that the SiNTs were characterized by

photo-electrochemical measurements that showed n-type conductivity and optical gap of

~1.6 eV. Quitoriano et al. [145] synthesized single-crystalline SiNTs using VLS along with

template preparation. Ge nanowires were first deposited using VLS method which acted as

template for Si shell. Ge-cores were removed by enabling exposure of the Ge core to H2SO4

and H2O2. These NTs resonate mechanically and achieved a quality factor of ∼1800. Similar

approach was adopted by Ben-Ishai et al. [146] to synthtesize SiNTs. They have in addition

differentially and selectively functionalized the inner and outer surfaces of SiNTs with

organic molecular layers containing different functional groups and

hydrophobicity/hydrophilicity chemical nature, via covalent binding, to give nanotubular

structures with dual chemical properties.

Other reports of synthesis

Other reports include the work of Tang et al. [147] in which they have prepared self-

assembled SiNTs with closed caps using one-dimensional silicon monoxide powder under

supercritically hydrothermal conditions (470°C, 6.8 MPa). The SiNTs had hollow inner

pore, crystalline silicon wall layers with a 0.31 nm inter-planar spacing and 2–3 nm

amorphous silica outer layers. Amorphous bamboo-like SiNTs were synthesized by a La


53

modified thermal evaporation process by Yuesheng Li et al. [148]. Qiu et al. [149] have

successfully fabricated SiNWs with undetached SiNTs by etching using aqueous HF and

AgNO3 solution. Xie et al. [150] have synthesized self-assembled SiNTs with diameters ~

50 - 80 nm and tubular wall thickness of ~ 10 - 15 nm using dual-RF-plasma treatment

technique and Cu catalysts. Saranin et al.[151] obtained highly ordered honeycomblike

nanostructure arrays using submonolayer Be deposition onto the Si(111) 7 X 7 surface held

at 500-700 C under ultrahigh vacuum conditions. The composition, structure, and

properties resemble those of Be-encapsulated Si nanotubes predicted by theory [152].

Applications of SiNTs

SiNTs have a great potential for photoemission applications due to quantum

confinement effects, so they can be seen as a part of future optoelectronics [153]. SiNTs also

show better electrochemical performance and hence they can be used in anodes of Li-ion

batteries [154–156]. It was the report of Park et al. that showed the capacity of SiNTs in Li-

ion fuel cell demonstrating a 10 times higher capacity than commercially available graphite

even after 200 cycles. This reported created interest in the scientific community about SiNTs

in Li-ion battery. The SiNTs here were prepared by reductive decomposition of a silicon

precursor in an alumina template and etching [155]. Furthermore, Yao et al. [157] studied

the field emission properties of SiNTs synthesized through a dry etching process in an ICP

system. Song et al. [158] using template method, synthesized arrays of sealed SiNTs and

used them as electrodes in lithium ion batteries. Jaehwan Ha and Ungyu Paik [154] used the

SiNTs synthesized by Song et al. [158] and modified process was used that showed

improved performance. Jung-Keun Yoo [159] used template method and fabricated SiNTs

through facile surface sol–gel reaction on easily obtainable organic nanowires and a simple

magnesium reduction. The fabricated SiNTs showed excellent electrochemical performances

when used as anodes for lithium rechargeable batteries. Zhenhai Wen et al. [160] used

template method involving preparation of silica nanotubes using rod-like NiN2H4 as a

template and the resulting silica nanotubes were then converted to Si nanotubes by a thermal

reduction process assisted with magnesium powder. The electrochemical properties of Si

nanotubes were investigated as anode of Li-ion batteries.


54

2.1.3.1 Thermal plasma assisted synthesis of SiNTs

Laser ablation induced thermal plasma has been successfully used by Yamada and

Fugiki [161] for growing multiwalled SiNTs composed of rolled quasi two dimensional

honeycomb structure with cylindrical symmetry as evidenced by high resolution TEM

analysis. Number of walls varied between 4 and 30; with outer diameter of 1 to 6 nm. The

interwall spacing was determined to be 0.36 nm and was thought to arise from puckered

structure of silicon atoms. Silicon two dimensional sheets are supposed to be rolled up about

different lattice axes to form tubes with different helicities (zigzag, armchair and some

helical tubes).

Polycrystalline SiNTs, filled with single crystal Sn have been synthesized by Feng et

al. [162] by using DC arc discharge method. Their Nanotubes have tapered structures with

homogeneous diameters of about several ten to several hundred nanometers. They have also

studied the PL properties of these NSs.

SiNTs grown by DC arc plasma assisted gas phase synthesis in our group have been

reported earlier [163,164]. TEM investigations revealed the presence of SiNTs and

nanoparticles in the ratio of 1:10. The tube diameter varied between 2 to 35 nm and they

were more than few hundred of nanometers long. Atomically resolved STM image showed a

honeycomb lattice feature in the centre showing an atomic arrangement compatible with

puckered Si (111) structure. Electron energy loss (EEL) measurements indicated that these

SiNTs were formed by single or very few silicon layers. Investigations related to chirality

were carried out by measuring the I-V curves using scanning tunneling spectroscopy for

several atomically resolved tubes. In-depth characterization using EEL near edge spectra

(EELNES) provided useful information about the non oxidized SiNTs and the presence of

oxide on Si nanoparticles.

2.2 Silicon carbide

Silicon carbide bears exceptional advantage because of its semiconducting properties

added with high thermal and chemical stability, and good hardness properties. Due to these

properties it finds applications in varied fields along with electronics industry. Thus, SiC

spans various fields of research. Nanotechnology has provided new avenues for research in

SiC.


55

2.2.1 Synthesis of SiC nanostructures

Acheson process

If we go through the history of SiC, it is found that it is older than our solar system,

having wandered through the Milky Way for billions of years as generated in the

atmospheres of carbon rich red giant stars and by supemova remnants. The possibility of Si

and C bond was first suggested by Jöns Jakob Berzelius in 1824 [165]. Years later, it was

accidently synthesized by Acheson in 1890 from mixture of coke and silica and he found

that it could substitute diamond as an abrasive and cutting material. This process remains

most used even today for industrial purposes and synthesis of SiCNPs is also carried out

using this method by modifying it [166]. Carbothermic reaction of spherical amorphous

SiO2 NPs (size of 2-10 nm) with sucrose at 1500 °C (with additional annealing at 700°C for

decarburization) results in preparation of 6H- and 4H-SiC polytypes with high density of

stacking faults and small amount of 2H and 3C SiCNPs with typical size of 5-10 nm [167].

The SiCNW preparation is possible by evaporation at 1600°C in argon atmosphere using

different initial products: silicon in graphite crucible and a graphite substrate[168] or a mix

of powders of silicon and graphite activated by grinding [169].

Other widely used synthesis techniques

The other widely used techniques include sol gel synthesis and mechanical alloying.

Sol-gel synthesis has several outstanding features such as high purity, high chemical activity

besides improvement of powder sinterability. Nevertheless, this process suffers due to large

duration of synthesis and high cost of the raw materials. This technique have been explored

by Julbe et al. [170] Raman et al. [171], Zheng et al. [172], Sharma et al. [173] and others.

Different morphologies of SiC nanostructures like nanowires [174], nanoparticles [175] etc.

have been synthesized by this method.

Mechanical alloying is a solid state process capable of obtaining nanocrystalline SiC

with very fine particles homogeneously distributed at room temperature and with a low cost.

This synthesis method have been used by Chaira et al. [176], Rajamani et al. [177],

Eskandarany et al. [178], Aberrazak & Abdellaoui [179], Ghosh and Pradhan [180], etc.

Microwave assisted synthesis have been reported by Satapathy et al. [181], Aguilar et al.

[182], Moshtaghioun et al. [183] and several others. A combinatorial approach using heating


56

and ball milling was adopted by Wang et al. to synthesize 60-100 nm diameter of SiC

nanowires [184].

Other synthesis techniques with a perspective of electronics

The electronic application of SiC was first observed in 1907, when H.J. Round

produced the first Light Emitting Diode (LED). However, SiC evolved as a semiconductor

material in 1955 only when Lely presented a new concept of growing high quality SiC

crystals. It was then used in LED followed by the power devices and high frequency

devices. But, like silicon, SiC is also an indirect band gap semiconductor, thus poor for

optoelectronic devices [185]. Secondly, miniaturization demands reduced dimension of SiC.

CVD is one of the suitably used methods to produce SiC in various forms; thin films [186],

nanoparticles [187,188], whiskers, nanowires [189] and nanorods [190]. Other techniques of

synthesis used for devices include electrochemical and chemical etching [191–193],

lithography [194], implantation of carbon ions in silicon [195] and joint implantation of ions

of carbon and silicon in SiO2 matrixes [196]. The carbidization study of porous silicon,

prepared by electrochemical etching, has revealed the 3C-SiC nanocrystal formation with

the size of 5-7 nm at temperature 1200-1300°C [197]. In a number of works (see, for

example, [195, 198–200]), the features of the SiC nanoparticle implantation (ion beam) are

investigated. Chen et al. [201] have used pyrolysis of silicone to synthesize SiC wiskers.

The as-prepared amorphous SiC particles were synthesized from the decomposition of

tetramethylsilane precursor in a plasma, operated at room temperature and low precursor

partial pressure (0.001-0.02 torr) using argon as carrier gas (3 t0orr). The synthesis

conditions were varied to prepare nanoparticles in the size 4-6 nm with reasonable

monodispersity.

Synthesis of SiC nanotubes (SiCNTs) and nanosheet

Synthesis of SiCNTs can be carried out by various methods, for example, interaction

of SiO vapor with carbon nanotubes [202] or reduction of methyltrichlorosilane by hydrogen

in the presence of the catalyst and co-catalyst (ferrocence and thiophene) at temperature

1000°C during 1 hour [196]. Outer diameter of nanotubes was about 20- 80 nm, the value of

inner diameter fluctuates in the range of 15-35 nm [196]. The detailed thermodynamic

analysis of the β-SiC gas-phase deposition using methyl-trichlorsilane as precursor, has been

realized in work [203] in which it was shown that optimum temperature of reduction by


57

hydrogen for the β-SiC preparation is about ~1200°C. Zou et al. have reported low-

temperature solvothermal route of synthesis of 2H–SiC nanoflakes [204]. Several

researchers have studied the possibility of two dimensional sheet-like [205–207] and CNT-

like structures and studied their properties. Lin et al. [208] have reported the synthesis of

light-emitting two-dimensional ultrathin silicon carbide.

Review articles and books on SiC nanostructures

There are several review articles on synthesis techniques like articles by R.A.

Andrievski [209,210]. The nanosized SiC synthesis methods as applied to electrical/optical

properties up to 2005 are discussed in a review by Fan et al. [198] with a comprehensive

description of technological regimes.

2.2.1.1 Thermal plasma assisted synthesis of SiC nanostructures

Thermal plasma route is one of well known route used for its synthesis which gives

the product in one step. In this type of synthesis silicon source and carbon source are

subjected to thermal plasma where the reaction between Si and C takes place to yield SiC.

Y. Leconte and co-authors [211] have synthesized SiC nanoparticles in ICP system of size

between 20 and 40 nm using SiC micron sized powder. They have used Ar-H2 during

synthesis due to which the decarburation of SiC took place and resulted in the formation of

some silicon impurities. To avoid silicon impurities excess of methane gas was passed after

which the formation of silicon was fully avoided. SiC nanopowder was synthesized on large

scale using SiC powder (with an average size about 1-3 µm) inside the dense plasma formed

by the ICP torch [212]. Sang-Min Ko and co authors [213] used organic sources; different

types of silane for SiC synthesis. They found Si as well as C impurities in the product

alongwith SiC nanoparticles. Carbon was removed by heat treatment and then silicon was

removed by treatment with HF.

A direct plasmodynamic technique was used by A. A. Sivkov [214] in which they

used solid silicon and carbon powder sources for the synthesis where they could obtain pure

beta silicon carbide particle with a wide size distribution between 10 - 500 nm alongwith

some silicon and carbon impurities. They have varied the ratio of silicon to carbon but kept

the proportion of silicon greater than carbon. Karoly et al. [215] have synthesized SiC

nanoparticles by RF thermal plasma method. Precursor mixtures comprised commercial

silica powder and various types of carbon source including graphite, charcoal, carbon black


58

as well as the carbonaceous residue of tyre pyrolysis. The obtained SiC consisted of

nanosized particles that were crystallized mainly in β phase with traces of α. The conversion

rate of the silica precursor to SiC varied between 60% and 73% depending on the type of

carbonaceous material and on the excess carbon.

Prabhakar Rai and co authors [216] have synthesized silicon carbide nanoparticles by

passing 1-5 µm sized particles through non-transferred thermal plasma. They were

converted in to the spherical particles of silicon carbide with β-SiC as the major product.

Some silicon impurity was also observed in the nanoparticles. Seung-Min Oh and co authors

[217] have synthesized SiC nanoparticles of average size 71 nm. They used precursors SiCl4

and CH4 for the synthesis. Initially when the synthesis was performed in presence of only

argon gas, large percentage of impurities was observed in the form of Si and graphite

particles. After addition of large percentage of H2 during synthesis the percentage of SiC

increased considerably with a small amount of SiO2 impurities.

B. B. Nayak and co-authors [218] have designed a plasma reactor to synthesize

silicon carbide nanoparticles from rice husk in a DC arc plasma reactor. The work done by

them is remarkable as the precursor used is a waste product. But, the particles obtained are

of micron size with lot of impurities present. B. B. Nayak and co-authors [219] in other

publication reported synthesis of Silicon carbide dendrite (micrometer size) by carbothermic

reduction of rice husk ash in an arc plasma. Transmission electron microscopy reveals the

occurrence of equispaced primary arms (60–70 nm in length) in the dendrite consisting of

nanorod bundles. Each nanorod is seen to contain thin transverse lamellas, which appear like

slip bands/twins in atomic layer thickness.

Jian Zhang and co authors [220] have synthesized silicon carbide nanowires by using

DC direct arc thermal plasma assisted vapour phase synthesis. They used a graphite

crucible as anode in which equimolar mixture of Si, SiO2 and C was placed. The 8 mm thick

tungsten cathode was used for arcing. The deposit obtained on the cathode surface were SiC

nanowires with diameter from 100-200 nm and 10-20 μm in length. The nanowires formed

consisted of core of SiC wraped with SiO2 layer.

Study of thin film deposition using thermal plasma have been carried out by

Girshick et al .[221] and Rao et al. [222]. They observed that the film in thermal plasma is

grown by direct nanoparticle impact. Thus, concluded that thermal plasma is more suitable


59

for particle synthesis. Peric et al. [223] have investigated the synthesis process of solid SiC

theoretically by computing the equilibrium composition of the gas mixtures involving

silicon and carbon in the presence of argon and hydrogen at various silicon/carbon amounts

and at two different total pressures in the system, in the temperature range between 1000 and

6000 K.

2.2.2 Applications of SiC nanostructures

Microwave applications

SiC is considered to be one of the important microwave absorbing materials due to

its good dielectric loss to microwave [224]. SiC can absorb electromagnetic energy and can

be heated easily with a loss factor of 1.71 at 2.45 GHz at room temperature [225, 226]. Ye et

al.prepared nano-SiC/N solid solution powders by laser method and studied the dielectric

properties at a frequency range of 8.2-12.4 GHz [227]. Li et al. [228] investigated the

microwave absorption properties of B-doped SiC powders synthesized by sol-gel process.

Wu et al. [229] investigated microwave absorbability of single-crystalline β-SiC nanowires

(diameters ~20–80 nm and lengths ~10μm) synthesized by a reaction of CNTs and silicon

vapor from molten salt medium at 1250°C by dispersing it in silicone matrix. Yang et al.

[230] have studied the temperature-dependent dielectric properties and enhanced microwave

absorption at gigahertz range (8.2–12.4GHz).

Electron field emission applications

Electron Field Emission application of nano-SiC have been widely studied, few

reports of which are mentioned here. Kang et al. [231] have studied the field emission from

nanoporous SiC. Meng et al. [232] used template assisted catalyst free CVD for synthesis of

SiC nanobelts and studied its field emission properties. Zhou et al. [233] studied the field

emission properties of β-silicon carbide nanorods (diameter ~ 5–20 nm; length, 1 μm),

grown on porous silicon substrates by CVD with an iron catalyst. Wu et al. [234] have

synthesized well-aligned SiC nanowire arrays on carbon cloth by a facile CVD and found

they are excellent field emitters. Chen et al. [235] have studied the field emission from

hexagonal prism-shaped single-crystal 3C-SiC nanowires with high aspect ratio grown on

graphite substrate. These NWs showed a very low threshold field of 2.1V μm−1

, high

brightness and stable field emission performance.


60

Fillers

SiC nanoparticles are used to increase the thermal and mechanical stability of

polymers e.g. 1 wt. % loading of silicon carbide (β-SiC) nanoparticles shows improvement

in both thermal and mechanical properties of SC-15 epoxy resin when compared to the neat

system [236]. Huang et al. [237] have studied the effects of a coupling agent and the content

of the SiC filler on the filler dispersion and the mechanical and thermal properties of the

polyethylene/nano silicon carbide composites.

Other applications

Besides, nano SiC have been studied for many other applications like Ivekovic et al.

[238] have provided an overview of the main characteristics of SiCf/SiC that suggest the use

of this SiC-based composite as a structural material for the blanket in future fusion reactors.

Hosoya et al. [239] have performed the polishing experiments on resins and found that the

SiC particles less than 12 μm possessed good polishing properties. Kriener et al. [240] found

that heavily boron doped 3C-SiC and 6h-SiC exhibit a similar critical temperature and field

strength and are type-I superconductors. Celata et al. [241] have studied the nanofluid

coolant application of SiCNPs. Silicon carbide is used to prepare graphene [242] and

diamond [243]. Pan et al. [244] have studied the field emission properties of SiC nanowire

arrays. Shafiei et al. [245] have investigated the application of hydrogen gas sensing

properties of SiC - graphene contacts.

SiC being biocompatible [246] in nature has many biomedical applications like

nano-SiC in bone implantation [247], as semipermeable membrane [248], bioanalytical

assays, cell imaging, biosensors [249], etc.

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70

Chapter 3

Experimental Techniques &

Procedures

This chapter presents the significant experimental techniques and methodology used in

carrying out the research work

Chapter 3. Experimental techniques and procedures

71

In this chapter the significant experimental techniques and methodology used in

carrying out the research work have been discussed. Initially, the plasma system used for the

syntheses of different nanostructures (NSs) and mechanism of synthesis is explained. In the

second part, aspects about some of characterization tools are described. The instrumentation

and the basic concepts behind every technique are not described. There is only an overview

of the perspectives which were helpful in this work. Only a few of the techniques are

explained with the emphasis on ‘How to analyze the data?’ especially about transmission

electron microscopy.

Details about the parameters used to synthesize the NSs and the experimental

conditions used to study them have been included in the chapters consisting of ‘Results and

Discussion’. The experimental details used to study the applications of the NSs have also

been included in the respective chapters discussing the results of synthesis of respective

nanostructures.

3.1 Experimental method of synthesis

DC direct arc plasma assisted gas phase condensation was used for synthesizing the

chosen nanomaterials. The experimental system is versatile and simple and has been earlier

used by several researchers in this group for synthesizing metals nanoparticles like silver

and various compounds like AlN [1], Al2O3 [2], LaB6 [3] apart from Carbon Nano Tubes [4]

and Graphene [5].

3.1.1 DC direct arc thermal plasma set up

The system necessarily consisted of an arcing assembly with an anode and cathode.

The arc in such systems is initiated simply by applying sufficient voltage between the two

electrodes while maintaining literally no gap between the two. Once the arcing begins the

electrodes are withdrawn so as to maintain a plume of plasma in the region. The schematic

of the DC direct arc thermal plasma reactor is show in figure 3.1 (a) and its photograph is

shown in figure 3.1 (b).

The reactor consisted of a water-cooled stainless steel (SS) cylindrical retort. It can

be divided in to two parts upper chamber and lower chamber. Lower chamber has a height

of 140 mm and an inner diameter of 240 mm, while upper chamber has a height of 300 mm

and an inner diameter of 250 mm. A movable SS hollow rod (outer diameter 32 mm) is


72

fitted at the centre of lower chamber through a Wilson seal. On the top of this SS hollow rod

a copper cup is fixed on which a graphite crucible is mounted. In this graphite crucible the

precursor, whose synthesis is to be done, is placed. This whole assembly is the anode.

Similarly, a movable hollow rod (diameter 2 mm) is fitted on the centre of upper chamber on

which a tungsten rod can be mounted that acts as a cathode. Tungsten rod of diameter 4 mm

is used, sometimes with a graphite cap. Arcing is done between the precursor and the

cathode. The geometry of the electrodes was modified for different syntheses. The

electrodes are connected to a DC power supply from Kejearc (Type KTC 200, Current 50-

200 A and open circuit voltage of 65 V, used for welding purposes).

Figure 3.1 (a) The schematic of the DC Direct arc plasma reactor used for the synthesis of silicon

and silicon carbide nanostructures, and (b) the photograph of the DC Direct arc plasma reactor.

Both the electrodes have a facilitation of water circulation. Water is circulated

throughout the doubled walled chambers and electrodes through a chiller that can provide

cooling between 288 K and 298 K. There is facilitation for two ports (on which CF48

(conflat flange 48) can be fitted) in the diametrically opposite directions of lower chamber.

One of these ports is used to connect a motor drive rotary pump (2L, 3 Phase) through a


73

CF48 to (Kwik flange) KF10 converter used for the evacuation of the chamber. The

evacuation is regulated by a butterfly valve. Perpendicular to these ports there is

arrangement for two CF28 flanges of which one is connected to the gas cylinder through

CF28 to KF10 converter used for the filling of chamber with the desired gas. The flow of

gas is regulated by a throttle valve.

Upper chamber has a view port mounted on an oblique column by CF48. The

cathode can be moved manually by a screw arrangement.

3.1.1.1 Electrode assembly for the synthesis of silicon nanostructures

The schematic of electrode assembly used for the synthesis of silicon nanoparticles

(SiNPs) is shown in figure 3.2 and the photographs are shown in figure 3.3.

Figure 3.2 The schematic of the electrode assembly used for the synthesis of silicon nanoparticles (a)

anode assembly showing SS hollow rod, copper cup and cylindrical graphite crucible (CR1) marked

by 1, 2 and 3 respectively, and (b) cathode consisting of tungsten rod marked by 4.

Figure 3.3 The photographs of the electrode assembly used for synthesis of silicon nanoparticles (a)

the anode assembly showing SS hollow rod, copper cup and cylindrical graphite crucible (CR2)

marked by 1, 2 and 3 respectively, and (b) cathode consisting of tungsten rod marked by 4.


74

The electrode assembly (Figure 3.2 & 3.3 (a)) used for the synthesis of SiNSs

consisted of graphite crucible (anode) of diameter 30 mm with a cylindrical cavity of depth

about 9 mm in which silicon microcrystalline powder is placed. This crucible will be further

referred as CR1. The cathode consists of a tungsten rod of diameter 4 mm tapered at the end

for arcing.

3.1.1.2 Electrode assembly for the synthesis of SiC nanostructures (SiCNSs)

Two different sets of crucibles were used; first with the graphite crucible (diameter

3cm) having a conical cavity (Figure 3.4 (a) and Figure 3.5 (a)) in which silicon precursor

was kept. The second set of crucible consisted of two stage graphite crucible; first stage

(diameter 30mm) indicated by 3 and second stage by 4a (diameter 30mm), 4b (diameter

17mm) and 4c (diameter 10mm) respectively in figure 3.4 (b). The crucibles in figure 3.4

and figure 3.5 will be further referred as CR2 (Figure 3.4 (a)), CR3 (Figure 3.4 (b), 4a), CR4

(Figure 3.4(b), 4b) and CR5 (Figure 3.4(b), 4c). The reason of using two stage crucible was

i) to completely cover the copper electrode with the graphite cup having diameter same as

that of the copper cup so as to avoid direct arcing between the cathode and copper cup, ii) to

cool copper cup but avoid cooling of the crucible so as to increase evaporation of the

precursor. Cathode consisted of a tungsten rod of diameter 4 mm on which a graphite cap of

diameter 9 mm was fitted.

Figure 3.4 The schematic of the electrode assembly used for the synthesis of silicon carbide

nanoparticles (a) the anode assembly showing SS hollow rod, copper cup and conical graphite

crucible marked by 1, 2 and 3 respectively, (b) the anode assembly showing SS hollow rod, copper

cup and first stage graphite crucible marked by 1, 2 and 3 respectively, 4a (CR3), 4b (CR4) and 4c

(CR5) represent second stages of crucibles, and (c) the cathode consisting of tungsten rod fitted with

a graphite cap marked by 5.


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Figure 3.5 The photograph of the electrode assembly used for the synthesis of silicon carbide

nanoparticles (a) the anode assembly showing SS hollow rod, copper cup and conical graphite

crucible marked by 1, 2 and 3 respectively, (b), (c) and (d) consist of anode assembly showing SS

hollow rod, copper cup and two stage graphite crucibles marked by 3 (first stage) and 4a(CR3), 4b

(CR4) and 4c (CR5), and (e) the cathode consisting of tungsten rod fitted with a graphite cap marked

by 5.

3.1.2 Synthesis procedure and mechanism of synthesis

Before synthesis, the chamber is evacuated using a rotary vacuum pump to a

pressure of 10-3

mbar. It is then filled with the desired gas and again evacuated. This is

repeated a few times in order to reduce gaseous impurities present inside the chamber. We

repeated this process thrice. The desired pressure is maintained throughout the experiment

with the help of a butterfly valve. For all the synthesis runs, the anode (lower electrode)

along with the chamber is grounded. The cathode (upper electrode) was biased negatively

with the help of a DC power supply. The electrodes are then approached for arcing. Once

the arc is generated they are retracted to form a stable plasma plume. The parameters of the

reactor are mentioned in Table 3.1.

Table 3.1 The parameters of DC arc Plasma reactor

Reactor parameters Range

Arc current 50 -200 A

Arc Voltage 5-60 V

Chamber pressure 100 -760 torr

Cooling Temperature 278-288 K

Due to the heat of plasma plume the precursor melts and eventually gets evaporated.

A steep temperature gradient exists at the edge of the plasma, which leads to super cooling,

nucleation and growth and hence to the formation of nanoparticles. The whole process of

growth induced by thermal plasma is schematically represented in Figure 3.6.


76

Figure 3.6 The schematic of the process of growth induced by thermal plasma [6].

3.2 Characterization techniques

The characterization tools used in this research work include X-Ray Diffraction

(XRD), UV-Visible (UV-Vis) Spectroscopy, Fourier Transform Infrared (FTIR)

Spectroscopy, Raman Spectroscopy, Scanning Electron Microscopy (SEM), Scanning

Tunneling Microscopy (STM), Transmission Electron Microscopy (TEM), High Resolution

Transmission Electron Microscopy (HRTEM), Electron Energy Loss Spectroscopy (EELS)

and Thermogravimetry (TG).

XRD was used to study the crystalline nature of nanoparticles. The crystalline phases

were determined by comparing the XRD line positions with the standard data using JCPDS

files. The best possible Bragg angles and the intensity ratios were used to draw the

conclusions about the chosen crystalline phase of a structure. The average size of the

crystallite was determined from the full width half maximum (FWHM) of the XRD line

using the Scherrer formula [7] given by,

, (3.1)

where, t is the mean size of the ordered (crystalline) domains, which may be smaller or

equal to the grain size; K is a dimensionless shape factor, with a value close to unity. The

shape factor has a typical value of about 0.9, but varies with the actual shape of the

crystallite; λ is the wavelength of the X-ray source; β is the line broadening at half the

maximum intensity (FWHM), after subtracting the instrumental line broadening, in radians;

θ is the Bragg angle.

http://en.wikipedia.org/wiki/X-ray


77

FTIR Spectroscopy was used to find out the presence of polar bonds in the

nanoparticles. In case of silicon and silicon carbide nanostructures it was especially the Si-O

and Si-H bonds that were detected. The FTIR absorption peaks corresponding to different

Si-O and Si-H modes are shown in Table 3.2.

Table 3.2 FTIR absorption peaks corresponding to different vibrations of bonds of Si with O2, H2 and

C [8,9]

Wavenumber Corresponding IR vibrations

450 – 470 cm-1

Si-O-Si rocking modes

600 – 660 cm-1

Dipole vibrations of various SiHx groups

790 cm-1

The stretching mode Si–C bond

812 cm-1

Si-O-Si bending modes

800 – 890 cm-1


915 cm-1

SiH2 scissor mode

930 – 950 cm-1

The stretching vibration of Si–OH

Around 1000 cm-1

Si-O vibrations

790 – 800 cm-1

, 1060 – 1090

cm-1

, 1160-1190 cm-1

Vibrational stretching of the Si–O–Si bond

2000 – 2200 cm-1


2100 cm-1

Vibration mode of hydrogen bonded to the surface of

crystalline Si

2250 cm-1

Vibration mode of hydrogen bonded to the surface of

amorphous Si

2125 and 2245 cm-1

The stretching mode of SiH2 group

UV-Vis Spectroscopy was used to study the optical properties of the nanoparticles.

It was also used to study the band gap of SiC-nanoparticles.

Raman spectroscopy is a powerful tool that can be used to determine the solid state

structure. The Raman spectrum can be used to determine the Si-Si bond thus the presence of

crystalline silicon. The crystalline silicon gives a symmetric Raman peak at 520.5 cm-1

with

an FWHM of 2.8 cm-1

. The amorphous silicon peak is totally distinct from this and can be

easily distinguished. Besides, due to reduced dimensions there is shift in the peaks of Raman

spectra. Each and every different bond in Si shows different signature. Thus, it is a powerful

tool to study silicon.


78

SEM is used to study the microstructure and particle surface morphologies. STM is

used to observe the tubular nature of silicon nanotubes and study the atomic arrangement in

them. TG analysis was used mainly for SiC nanoparticles to find the presence of silicon and

carbon impurities in the as synthesized samples.

3.2.1 Transmission electron microscopy

TEM has been used extensively in this work. If required to describe TEM in a very

simple language, it can be described as a tool that captures the shadows of the objects in the

path of electron beam. Consider a light source with obstacles between source and screen, the

shadow casted on the screen would depend on the nature of obstacle like shape, thickness

and opacity of the obstacles. Similar phenomenon is observed in TEM. Difference is that

optical source is replaced with high energy (~100 keV) electron beam and the obstacle

consists of nanoparticles that are to be analyzed. And since it is a high energy electron beam

there are two different types of interactions that take place between electron beam and the

sample particles.

Elastic Interactions: No energy is transferred from the electron to the sample. The electron

either passes without any interaction (direct beam) or is scattered by the electrostatic

attraction of the positive potential inside the electron cloud. The arising signals are mainly

exploited in imaging and electron diffraction in TEM.

Inelastic Interactions: Energy is transferred from the incident electrons to the sample

producing secondary electrons, phonons, UV quanta or cathodo-luminescence. Ionization of

atoms by removing inner shell electrons leads to the emission of X-rays or Auger electrons.

These signals are used in analytical electron microscopy.

3.2.1.1 Different modes of TEM imaging

TEM is a technique that uses the interaction of energetic electrons with the sample

and provides morphological, compositional and crystallographic information. The electron

emitted from the filament passes through the multiple electromagnetic lenses and make

contact with the screen where the electrons are converted into light and an image is

obtained. The speed of electrons is directly related with the electron wavelength and

determines the image resolution. A modern TEM is composed of an illumination system,

condenser lens system, an objective lens system, magnification system, and the data

recording system. A set of condenser lens focus the beam on the sample and an objective

http://www.microscopy.ethz.ch/elast.htm


79

lens collects all the electrons after interacting with the sample and form image of the sample,

and determine the limit of image resolution. Finally, a set of intermediate lenses magnify

this image and projects them on a phosphorous screen or a charge coupled device (CCD)

camera.

I. Modes using elastic scattering

a. Bright Field Imaging

In the bright field mode of the TEM, only the direct beam is allowed to contribute in

image formation. Scattered electrons are efficiently blocked by the aperture. It is essentially

the weakening of the direct beam’s intensity that is detected by the bright field imaging. A

main component of this weakening is the mass-thickness contrast or diffraction contrast.

This contrast can be explained by simple model of elastic scattering by Coulomb interaction

of electrons with the atoms in a material. Heavier elements represent more powerful

scattering centers than light element owing to the larger number of charges that the atom

carries. Due to this increase of the Coulomb force with increasing atomic number, the

contrast of areas in which heavy atoms are localized will appear darker than of such

comprising light atoms. This effect is the mass contrast. Furthermore, more electrons are

scattered in thick samples than in thin ones as the numbers of atoms, thus the scattering

centre that are lying in the path of the electron are larger. Therefore, thick areas appear

darker than thin areas of the same material. This effect leads to thickness contrast. Together,

these two effects are called mass-thickness contrast or diffraction contrast (Figure3.7). This

contrast is useful to observe size and shape of nanoparticles.

Figure 3.7 Schematic representation of contrast generation depending on the mass and the thickness

of a certain area [10].


80

To obtain lattice images, a large objective aperture has to be selected that allows

many beams including the direct beam to pass. The image is formed by the interference of

the diffracted beams with the direct beam (phase contrast). If the point resolution of the

microscope is sufficiently high and a suitable crystalline sample oriented along a zone axis,

then high-resolution TEM (HRTEM) images are obtained. The distance between the planes

can be directly calculated or the Fast Fourier transform (FFT) of the image can be obtained

which can be used to calculate the lattice spacing [10].

b. Dark field imaging

In dark field imaging the direct beam is excluded. If a sample is crystalline, then

another type of contrast appears in TEM images, namely diffraction or Bragg contrast. If a

crystal is oriented close to a zone axis, many electrons are strongly scattered to contribute to

the reflections in the diffraction pattern. A dark field image shows almost equally bright

image of the crystals that are responsible for the same diffraction. Thus, in dark field image,

the crystallites diffracting into a particular area of reciprocal space appear with bright

contrast whereas others remain less bright.

In bright field imaging the direct beam is allowed to pass thus the image obtained

shows particles in dark and screen brighter as the intensity of electrons is greater where there

are no obstacles. While in dark field imaging only the diffracted beam is allowed to form the

image, thus only the parts which diffract the beam appear brighter while the rest appear dark

as there are no electrons. Figure 3.8 shows the ray diagram of how the bright field and dark

field imaging is performed by using the different apertures.


81

Figure 3.8 (a) Left: bright-field mode, and (b) Right: dark-field mode [11].

c. Selective area electron diffraction (SAED)

Depending on the size of the investigated crystallites, different types of electron

diffraction patterns are observed. If exclusively a single crystal contributes to the diffraction

pattern, then reflections appear on well-defined sites of reciprocal space that are

characteristic for the crystal structure and its lattice parameters. Each set of parallel lattice

planes that occur in the investigated crystal and in the selected zone axis give rise to two

spots with a distance that is in reciprocal relation to that in real space (The normal to the

plane of the spot pattern is termed as the “zone axis”. The zone axis is also normal to the

electron beam). Thus, large d-values cause a set of points with a narrow distance in the

diffraction pattern, whereas small d-values cause large distances. If more than one crystal of

a phase contributes to the diffraction pattern, as it is the case for polycrystalline samples,

then the diffraction pattern of all crystals are superimposed. Since the d-values are same, the

corresponding distances in reciprocal space are same, the spots are then located on rings.

Such ring patterns are characteristic for polycrystalline samples [10]. Figure 3.9 shows the

ray diagram to obtain selective area electron diffraction pattern. Here, the aperture is used at

the back focal plane and the diffracted beam is allowed to pass and form the image.


82

Figure 3.9 (a) Left: ray diagram to obtain selective area diffraction pattern in TEM, and (b)

Right: geometry for electron diffraction and definition of camera-length, L. The electron wavelength

is λ, and the camera constant of (eqn 3.6) is λL [11].

Figure 3.9 (b) shows the geometry of path of electrons during diffraction. The

separation of the diffraction spots can be used to determine interplanar spacings in crystals.

Consider the geometry of a selected area diffraction pattern in figure 3.9 (b), which shows

the “camera-length,” L that is characteristic of the optics of the microscope. Bragg’s law is

given by,

2d sinθ = λ, (3.2)

where, d is the lattice spacing, θ is the angle of diffraction and λ is the wavelength of the

source.

Now θ ~ 1° for low order diffractions of 100 keV electrons (λ = 0.037 Å) from many

materials. For such small angles,

(3.3)

By the geometry of figure 3.9 (b),

(3.4)

Thus, from equations (3.2), (3.3) and (3.4),


83

rd = λL (3.5)

Thus,

(3.6)

Equation (3.6) is the “camera equation”. It allows us to determine an interplanar

spacing, d, by measuring the separation of diffraction spots, r. To do so, we need to know

the product, λL, known as the “camera constant”. In modern TEMs it is already calculated in

the TEM software and displayed in the form of calibrated scale (unit = nm-1

) [11].

SAED patterns from a single crystalline particle depend on the orientation of the

particle with respect to the zone axis. After obtaining the spot pattern the first thing to be

done is calculating the d-spacing (Lattice spacing) corresponding to the spots. From the d-

spacing the corresponding planes belonging to the material can be found by matching the

data with JCPDS data cards for the particular systems. But sometimes in case of allotropes

of same materials same d-spacing belong to more than one crystalline form. So it becomes a

problem to identify the exact phase. So, it should be taken into notice that every crystalline

phase gives peculiar spot pattern for a particular zone axis. The pair of diffraction spots is

obtained from the planes which are perpendicular to the zone axis. So, depending on the

crystal structure and symmetry number of diffraction spots will be observed. The angle

formed by the two adjacent spots at the centre spot must be equal to the angle between the

two planes in real lattice only then the particular crystal structure should be confirmed.

For example, here the diffraction patterns for silicon with FCC cubic structure have

been included. The diffraction patterns with different zone axes obtained by using software

Carine Crystallography 3.1 have been shown in figure 3.10 (a) – (g). The corresponding

orientation of the crystal with respect to beam is shown in the centre with lattice parameters

a, b and c respectively. When the zone axis is (100) the diffraction pattern obtained consists

of four diffraction spots making an angle of 90° with each other (Figure 3.10 (a)). Here, the

lowest order planes that are perpendicular to zone axis are (100) are (022). Hence two pairs

of diffraction spots are observed. Similarly, different symmetries occur for different zone

axes which give different diffraction patterns.


84

Figure 3.10 The diffraction pattern for FCC Si crystal obtained using software Carine

Crystallography 3.1 oriented in different directions: (a) (100) Zone axis, (b) (101) Zone axis, (c)

(111) Zone axis, (d) (211) Zone axis (e) (311) Zone axis, and (f) (331) Zone axis.

The diffraction pattern observed for hexagonal lattice with ABAB stacking sequence

obtained using software Carine Crystallography 3.1, for different zone axes have been

shown in figure 3.11 (a) – (d). The corresponding orientation of the crystal with respect to

beam is shown in the centre with lattice parameters a, b and c respectively.


85

Figure 3.11 The diffraction pattern for hexagonal lattice with ABAB stacking sequence obtained

using software Carine Crystallography 3.1 oriented in different directions: (a) (001) Zone axis, (b)

(101) Zone axis, (c) (110) Zone axis, and (d) (100) Zone axis.

II. Modes using inelastic scattering

Inelastic scattering of electrons is used in “Analytical TEM”, which uses two types

of spectrometry to obtain chemical information from electronic excitations.


86

a. Energy-dispersive X-Ray spectrometry (EDS)

In energy-dispersive X-ray spectrometry (EDS), an X-ray spectrum is acquired from

small regions of the specimen illuminated with a focused electron beam, usually using a

solid-state detector. Characteristic X-rays from the chemical elements are used to determine

the concentrations of different elements in the specimen.

b. Electron energy-loss spectrometry (EELS)

In electron energy-loss spectroscopy, we deal directly with the primary process of

electron excitation, which results in the fast electron losing a characteristic amount of energy

due to inelastic interactions of the electrons. The transmitted electron beam is directed into a

high-resolution electron spectrometer that separates the electrons according to their kinetic

energy and produces an electron energy-loss spectrum showing the number of electrons

(scattered intensity) as a function of their decrease in kinetic energy. Information about the

internal structure can be obtained from this spectrum. For 100-keV incident energy, the

specimen must be less than 1 μm thick and preferably below 100 nm. The basic details of

this technique have been elaborated in the further section.

3.2.2 Electron energy-loss spectrometry in TEM

The energy losses described above are very small (few eV to 1000 eV) as compared

to the energy of incident electron beam. Thus, an electron energy analyzer needs to be

highly resolved. The energy resolution of an energy analysis system is limited by energy

spread (ΔEs) in the electron beam incident on the specimen. Thus, a monochromator is also

equally important to energy analyzer.

Monochromator: Recent commercial microscopes have energy resolutions better than 0.1

eV. This is accomplished by starting with a field emission gun, often a Schottky effect gun,

followed by an electron monochromator, often a Wien filter. Wien filter consists of a region

with crossed electric and magnetic fields that are tuned to cancel for electrons of one

velocity only, which avoid deflection and pass through the exit aperture of the filter [12].

Energy analyzer: EELS spectrometer is mounted after the projector lenses of a TEM. The

heart of a transmission EELS spectrometer is a magnetic sector, which serves as a prism to

disperse electrons by energy. Since the energy losses are small in comparison to the incident

energy of the electrons, the energy dispersion at the focal plane of typical magnetic sectors

is only a few microns per eV. Electrons that lose energy to the sample move more slowly


87

through the magnetic sector, and are bent further upwards. A slit is placed at the focal plane

of the magnetic sector, and a scintillation counter is mounted after the slit. Intensity is

recorded only from those electrons that bent through the correct angle to pass through the

slit. A range of energy losses is scanned by varying the magnetic field in the spectrometer.

3.2.2.1 Characteristic features of electron energy-loss spectrometry (EELS)

Figure 3.12 shows an electron energy loss spectrum of an iron fluoride film. This is

included as an illustration.

Figure 3.12 Energy-loss spectrum of an iron fluoride film: (a) low-loss region with a logarithmic

intensity scale, and (b) part of the core-loss region, with linear vertical scale [13].

In general electron energy loss spectra can be divided into three regions as stated below:

Zero loss peaks: The first peak, the most intense for a very thin specimen, occurs at 0 eV

and is therefore called the zero loss peak (Figure 3.12). It represents electrons that did not

undergo inelastic scattering (interaction with the electrons of the specimen) but which might

have been scattered elastically (through interaction with atomic nuclei) with an energy loss

too small to measure. The width of the zero loss peak, typically 0.2–2 eV, reflects mainly

the energy distribution of the electron source.

Low loss region: This region includes the energy losses between the zero loss peak and

about 100 eV. Here, the plasmon peaks are the predominant feature. The information about

the sample thickness, local chemistry and structure is obtained from features in EELS

spectra caused by plasmon excitations and core electron excitations. The more intense the

plasmon peaks are, the thicker the investigated sample area is.

The low-loss features arise from inelastic scattering by conduction or valence

electrons. For example in figure 3.12 (a) the most prominent peak, centered around 22 eV

results from a plasma resonance of the valence electrons. The increase in intensity around 54


88

eV represents inelastic scattering from inner-shell electrons, in this case the M2 and M3

subshells (3p1/2

and 3p3/2

electrons) of iron atoms. Its characteristic shape, a rapid rise

followed by a more gradual fall, is termed an ionization edge; it is the exact equivalent of an

absorption edge in X-Ray Absorption Spectroscopy.

As the electron moves through the solid, the backward attractive force of the positive

correlation hole results in energy loss. The process can be viewed in terms of the creation of

pseudoparticles known as plasmons, each of which carries a quantum of energy equal to Ep =

hfp = (h/2π)ωp. A Plasmon (or a Plasmon oscillation) may be described as an oscillation of

the conduction electrons with respect to the positive ion cores of the crystal lattice with

frequency ωp. A Plasmon must be considered as an elementary excitation, i.e.a quasiparticle,

characterized by its energy and momentum. This leads to the classification of two types of

plasmons:

i. A bulk Plasmon which has momentum with a component normal to the surface.

ii. A surface Plasmon which has no component of momentum normal to the surface.

The relation between surface Plasmon ( ) and bulk Plasmon is given by [14],

, (3.7)

where, is the permittivity of the medium.

The surface excitations dominate only in very thin (<20 nm) samples or small

particles. In the case of a metal, bulk plasmons are not excited and surface excitations can be

studied alone. Free electron metals such as aluminum have sharper plasmon peaks than do

alloys of transition metals, which have a high density of states at the Fermi energy. For

semiconductors, Rafferty and Brown [15] pointed out that the low-loss fine structure

represents a joint density of states multiplied by a matrix element that differs in the case of

direct and indirect transitions. Assuming no excitonic states, their analysis showed that the

onset of energy-loss intensity is proportional to (E − Eg)1/2

for a direct gap and (E − Eg)3/2

for

an indirect gap [12].

If the energy-loss spectrum is recorded from a sufficiently thin region of the

specimen, each spectral feature corresponds to a different excitation process. In thicker

samples, there is a substantial probability that a transmitted electron will be inelastically

scattered more than once, giving a total energy loss equal to the sum of the individual losses.


89

In the case of plasmon scattering, the result is a series of peaks at multiples of the plasmon

energy. This can be observed in case of silicon as shown in figure 3.13.

Figure 3.13 The energy-loss spectra recorded from silicon specimens of two different thicknesses.

The thin sample gives a strong zero-loss peak and a weak first-plasmon peak; the thicker sample

provides plural scattering peaks at multiples of the plasmon energy [12]

The loss measurements are very useful for silicon samples as though the mean free paths

for Si and SiO2 are very similar, but the plasmon energies differ substantially: 17 eV in Si

compared to 23 eV in SiO2 [16].

High loss region: At an energy loss of 100 eV to 1000 eV, the signal intensity drops

rapidly. The ionization edges occurring at a higher energy loss in figure 3.10 arises at

fluorine K-edge (excitation of 1s electrons) followed by iron L3 and L2 edges (representing

excitation of Fe 2p3/2

and 2p1/2

electrons). The continuous background comes from electrons

that generate unspecific signals, most importantly the Bremsstrahlung radiation. As in an X-

ray spectrum, there are additional peaks at well-defined sites in the EELS above the

background. These ionization edges appear at electron energy losses that are again typical

for a specific element and thus qualitative analysis of a material is possible by EELS. The

onset of such an ionization edge corresponds to the threshold energy that is necessary to

promote an inner shell electron from its energetically favored ground level to the lowest

unoccupied level. This energy is specific for a certain shell and for a certain element. Above

this threshold energy, all energy losses are possible since an electron transferred to the

vacuum might carry any amount of additional energy. If the atom has a well-structured

density of states (DOS) around the Fermi level, not all transitions are equally likely. This

gives rise to a fine structure of the area close to the edge that reflects the DOS and gives

information about the bonding state. This method is called electron energy loss near edge


90

structure (ELNES). From a careful evaluation of the fine structure farther away from the

edge, information about coordination and interatomic distances are obtainable (extended

energy loss fine structure, EXELFS) [10].

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Chapter 4

Synthesis of Silicon Nanostructures

& their Applications

This chapter provides the experimental outcomes related to the effects of different synthesis

parameters on the morphology of silicon nanostructures. The study of the nanostructures has

been discussed using different characterization techniques. Later, investigations about the

applications of silicon nanotubes have been elaborated.

Chapter 4. Synthesis of silicon nanostructures and applications

92

The synthesis of silicon nanostructures (SiNSs) is important from the point of view

of basic studies as well as their applications. The importance of SiNSs, their properties and

applications are elaborated in chapter 1 and chapter 2. This chapter presents the

experimental results of the attempts to synthesize SiNSs and investigating the effects of

different parameters of synthesis on the morphology of NSs. First part of this chapter

focuses on the optimization of parameters for the synthesis of silicon nanotubes (SiNTs)

using argon and hydrogen as the plasmagen gases and on the possible mechanism of

formation of NTs in the optimized parameters. Further, in depth analysis has been presented

by changing the concentration of hydrogen on the morphology of SiNSs.

Second part of the chapter includes the applications of the synthesized SiNTs, under

the optimized parameters, in two different areas. The first is a report related to the

antimicrobial applications and the second relates to the study of its field emission properties

4.1 Synthesis and characterization of silicon nanotubes

4.1.1 Experimental details

This set of experiments primarily aimed to synthesize SiNTs. The synthesis was

undertaken by following the procedure described in section 3.1.2. A graphite crucible of

diameter 3 cm with cylindrical cavity as described in section 3.1.1.1 was used for placing

the precursor i.e. microcrystalline silicon powder (Sigma Aldrich, 98.5 % purity). The

silicon powder in a graphite crucible formed the anode while a tungsten electrode (diameter

4mm, described in section 3.1.1.1) acted as a cathode. The ambient pressure in the chamber,

during synthesis, was kept constant at 500 torr that was guided by the earlier work on silicon

in the lab. The voltage drop across the arc was maintained between 12 and 16 V as the

preliminary experiments showed that the arc remained stable for this voltage. Keeping the

pressure and the arc voltage constant throughout the entire syntheses, the arc current was

only varied. Initially, the experiments were performed in the presence of argon following the

reports published by the previous workers in the literature [1–3]. The results of

characterization showed the presence of particles and wire like growth that showed some

changes in the morphology which is further elaborated in this chapter. The tubular growth

was not observed. Thus, based on the reactions favoured by the reported theoretical works


*Due acknowledgement to Dr. N.P. Lalla and **Dr. V. Sathe 93

***Due acknowledgement to Sujoy Karan and Prof. Richard Berndt

[4], H2 was added to Ar and syntheses were performed at different arc currents and results

were studied. Details of the parameters used in different synthesis experiments performed

are listed in Table 4.1.

Table 4.1 Details of the parameters used during optimization of parameters for the synthesis of

SiNTs

Ambient gas Sample Name Arc current

Ar

Si1 80 A

Si2 100 A

Si3 120 A

Si4 140 A

Ar:H2

(95:5 %Mol)

Si5 80 A

Si6 100 A

Si7 120 A

Si8 140 A

The samples were characterized using X-Ray Diffraction (XRD) and Transmission

Electron Microscopy (TEM). XRD patterns of the samples were recorded with Bruker D8

XRD machine with CuKα radiations, Ni filter and graphite monochromator. TEM images

were recorded by Technai G2 twin TEM with a 200 keV LaB6 thermionic emitter and a

Charged Couple Device (CCD) camera. For recording TEM images, the samples were first

dispersed in isopropyl alchohol by sonicating in an ultrasonic bath. Two to four drops of

these dispersions were then poured on a holy carbon coated copper grid (mesh 200). These

measurements were carried out at Department of Physics, SPPU, Pune. Some TEM

measurements were carried out at UGC-DAE Consortium for Scientific Research, Indore*.

The TEM instrument at UGC-DAE Consortium for Scientific Research, Indore is from

Technai with 200 keV LaB6 fillament and CCD Camera. Raman spectra were recorded in a

backscattering geometry at room temperature, using a Jobin-Yvon Labram HR800

spectrometer with a He–Ne laser (λ = 632.81 nm) for excitation**.

Further, the NTs were investigated by OMICRON Scanning Tunneling Microscope

(STM) in ultra high vacuum (UHV) at room temperature. These measurements were carried

out at Institut für Experimentelle und Angewandte Physik, Christian-Albrechts-Universität

zu Kiel, D-24098 Kiel, Germany***.


*Due acknowledgement to Dr. Paola Castrucci and Prof Maurizio De Crescenzy 94

Also, the NTs were characterized by Energy Filtered - High Resolution Transmission

Electron Microscopy (EF-HRTEM) and Nanobeam Electron Spectroscopy (NES) and

Diffraction (NED). Energy Filtered - High Resolution Transmission Electron Microscopy

and Nanobeam Electron Spectroscopy and Diffraction were performed using a FEI TECNAI

12 G2 Twin (120 eV) apparatus equipped with a “post-column” GATAN GIF energy filter

and a 794 IF Peltier cooled slow-scan charge-coupled device multiscan camera. A droplet of

the raw synthesis product diluted in ethylene was used to disperse the NSs on a gold grid

(mesh 1000). The use of the gold grid allowed us to perform measurements on free-standing

NTs. To enhance image contrast and resolution, chromatic aberrations were reduced by

collecting only elastic electrons (E = 0). All the high resolution images were collected at

the “Scherzer defocus”, so that to optimize the transfer function of the optical system

balancing the effect of spherical aberration (Cs = 2.2 mm) against a particular negative value

of f (about 103 nm). In this case the image contrast represents the two dimensional

projection of the crystal potential [5]. All the NES and NED experiments were performed on

individual NSs by using a nanometer-sized (up to about 20 nm) coherent parallel electron

beam [6]. Actually, these joint measurements allowed us to directly visualize the area of the

specimen (about 50 nm in diameter) from which spectra and diffractions were acquired. It is

worth to underline that imaging the sample before and after the spectrum acquisition is of

fundamental importance to confirm that no electron beam damage or specimen displacement

occurred during the measurement. All these measurements were carried out and analyzed at

Dipartimento di Fisica, Università Roma Tor Vergata and Unità CNISM, via della Ricerca

Scientifica 1, 00133 Roma, Italy and Dipartimento di Tecnologie e Salute, Istituto Superiore

di Sanità, 00161 Roma, Italy*.

4.1.2 Results and discussion

A) Samples synthesized in presence of argon

4.1.2.1 XRD analysis

Although, XRD analysis was not very useful in order to observe the formation of

NTs, it was carried out to observe the phase formation of the as synthesized samples. Figure

4.1 shows the XRD pattern of standard JCPDS Card no. 050565 corresponding to diamond

Si, samples Si1, Si2, Si3 and Si4. It can be observed that each of the as synthesized


95

samples consisted of XRD peaks corresponding to Si(111), Si(220) and Si(311) planes. The

average crystallite size was not calculated from the Scherrer formula as the peaks observed

here appear to be formed of overlapping peaks of low intensity and high broadening and

other with higher intensity and lower broadening. From XRD pattern, it could be observed

that the crystallinity of the samples increased with increasing arc current.

Figure 4.1 X-Ray diffraction patterns of Si samples synthesized in ambient argon

4.1.2.2 TEM analysis

Figure 4.2 shows the TEM micrographs for samples Si1, Si2, Si3 and Si4. The insets

show the selective area electron diffraction (SAED) patterns. All the samples consisted of

similar morphologies showing spherical particles-like and wire-like structures. Sample Si1

showed the presence of spherical particles below 20 nm, few NWs with diameters less than

6 nm and some 15-20 nm diameters were seen to be formed by joining of particles. As the

current was increased to 100 A for Si2, the percentage of NWs increased as compared to

NPs. Also, the average diameter of wires increased to 10 to 15 nm. Sample Si3 consisted of

NSs similar to Sample Si2. In sample Si4, bigger, well defined spherical NPs near to 50 nm

and below, and long NWs with greater diameter than other samples were observed.

However, no signature of NTs was observed in any of these samples.


96

Figure 4.2 TEM micrographs of as synthesized Si samples in argon (a) Si1, (b) Si2, (c) Si3, and (d)

Si4 (Insets show the selective area electron diffraction pattern of the corresponding samples)

B) Samples synthesized in presence of argon and hydrogen (95:5 mole%)

The study of samples synthesized in presence of Ar revealed the presence of NWs.

So, further experiments were synthesized by adding H2. The tubular structure in Si is a

metastable state, so it would require higher energy for its formation. In the earlier reports

NTs were synthesized in presence of Ar alone [1–3] which might have arised due to

differences in the purity of the ambient at microscopic levels. Enthalpy of plasma can be

increased by increasing the current and voltage of arc, which was already performed and the

results were not encouraging. Therefore, H2 was added to increase the enthalpy of plasma

[7]. Secondly, hydrogen forms bonds with Si, capping it, and thus restricting its growth.

Apart from this, Seifert et al. [8] have suggested that the glow discharge of monosilane can

be a possible way of SiNT synthesis. Monosilane consists of Si and H2, and therefore Si and

H2, as precursors, might possibly yield SiNTs. Thus 5% H2 was added to Ar during

synthesis. The results obtained are discussed below.


97


Figure 4.3 shows the XRD pattern of standard JCPDS Card no. 050565

corresponding to diamond Si samples Si5, Si6, Si7 and Si8. It can be observed that peaks

corresponding to Si can be seen in all the samples.

Figure 4.3 X-Ray Diffraction patterns of as synthesized Si samples in presence of argon and

hydrogen in the ratio (95:5)

Sample Si5 which was synthesized at 80 A, showed a very small peak corresponding

to Si(111). The average crystallite size calculated from the FWHM of this peak using

Scherrer formula gave the value of 16 nm. This value seems to be arising from very few

bigger particles in the samples. So, the intensity of the peak is extremely low and the value

of average crystallite size is quite large. Sample Si6 and Si7 showed nearly similar pattern of

XRD. The average crystallite sizes calculated using Scherrer formula for samples Si6, Si7

and Si8 are 15 nm, 9 nm and 17 nm respectively. From the XRD patterns, it could be

observed that even in case of Ar-H2 mixture the crystallinity of particles increases with

increasing arc current.


Figure 4.4 shows the TEM micrograph of sample Si5 synthesized at 80 A, right inset

shows the magnified image of the tip of NTs whereas the left inset shows SAED pattern of

the sample. This sample consisted of many one dimensional structures as compared to all the

other samples. These one dimensional structures, differed from other samples in two ways

(i) the diameter varied between 9 nm and 30 nm and (ii) the tip of the structures was


98

different showing the open end as shown in the right inset of figure 4.4. Thus, these

elongated structures were NTs. These NTs appeared in bundles wound around each other

along with some NPs of silicon. The NTs and NPs were seen to be in the ratio of ~70:30.

The NPs are spherical in shape with size variation between 5-25 nm while the diameters of

the NTs range between 9 nm and 30 nm. A large number of tubes have the diameter of 14 ±

2 nm while the lengths are of the order of several hundreds of nm.

Figure 4.4 TEM micrograph of sample Si5, right inset shows the magnified image of a nanotubes

and left inset shows the corresponding SAED pattern of the nanotubes and nanoparticles.

Tubular formation is apparent from the circular open tip of a single NT as shown in

the right inset of figure 4.4. The inner part of the NT appears homogeneously bright whereas

the hollow opening is clearly visible at the tip. The thickness of the annular dark wall seen at

the tip appeared to be less than 1 nm. The dark lines running along the lengths, at the two

edges of the NTs, appear to have similar thickness. Further, the tubular nature was

confirmed and studied using STM, NES and NED which will be discussed in the next

subsection.

When the current was increased further, the tubular NSs were not observed. Figure

4.5 (a) showed the TEM micrograph of sample Si6 where the inset shows the corresponding

SAED pattern. This sample consisted of one dimensional elongated as well as spherical NSs.

Figure 4.5 (b) shows the magnified image of an elongated structure. Here, the tip does not

show the hollow structure as observed in previous sample. Thus, this sample consists of

NWs and NPs. The SAED pattern shows the presence of rings corresponding to Si(111),

Si(220) and Si(311) planes.


99

Figure 4.5 (a) TEM micrograph of sample Si6 where the inset shows the corresponding SAED

pattern, and (b) the magnified image of the tip of an elongated structure.

Figure 4.6 shows the micrograph of sample Si7 and the inset shows the

corresponding SAED pattern. This sample shows the increase in the formation of particles

as compared to earlier samples. Figure 4.7 (a) and (b) shows the TEM micrograph of sample

Si8. This sample mostly consisted of spherical NPs of silicon and few NWs. The average

particle size is found to be around 30 nm. Thus, with increasing current, the particle growth

is seen to increase as compared to the one dimensional growth of NSs.

Figure 4.6 TEM micrograph of sample Si7 and the inset shows the corresponding SAED pattern


100

Figure 4.7 (a) and (b) TEM micrographs of sample Si8 and the inset in (a) shows the corresponding

SAED pattern

Thus, SiNTs were formed in presence of Ar-H2 (95:5 molar ratio) at 80 A arc

current. Further, the SiNTs were studied using different characterization tools.

4.1.2.5 STM analysis of SiNTs (Si5)

Figure 4.8 shows the STM image of a single SiNT adsorbed on freshly cleaved

highly oriented pyrolytic graphite surface (HOPG). Almost every NT found during the

experiment, exhibited quite small topographic height compared to its diameter as displayed

here in the line profile (Figure 4.8 (b)) taken along the marked line in figure 4.8 (a).

Figure 4.8 (a) STM image (360 nm X 360 nm) of single silicon nanotube on HOPG; Vbias = 1 V,

Itunn= 0.95 nA, and (b) line profile along the yellow line drawn in (a)

Such a discrepancy occurs mainly from the geometrical convolution between the

STM tip shape and the NTs [9,10]. However, the dependence of the tunneling gap [11,12]


101

on the local conductivity should also be considered. For example, a fullerene C60 appears

almost 40% lower than its actual height in the STM topography even with a single atom at

the apex of the STM tip. Thus, it is not in general easy to obtain the real geometry of the

tube adsorbed on substrate. From the line profile, we estimated [13] the tube diameter, with

some errors, of approximately 11 nm while the lower limit of the estimated height is 1.4 nm.

This indicates a strong radial compression of the NT induced by van der Waals interaction

between the tube and the substrate. Similar kind of radial compression occurs for carbon

nanotube (CNT) in case of low number of inner shells or a single wall [14] contrary to

multiwall tubes. So, possibly the investigated tubes are few layered or single walled

structures.

4.1.2.6 Raman analysis of SiNTs (Si5)

The presence of crystalline silicon was also confirmed using Raman spectroscopy.

Figure 4.9 (a) and (b) show the Raman spectra for crystalline silicon and sample Si5

respectively.

Figure 4.9 Raman Spectra of silicon samples (a) crystalline silicon, and (b) sample Si5.

The peak for crystalline silicon is found to be symmetric and appears at 520.5 cm-1

with full width at half maximum (FWHM) of peak to be 2.8 cm-1

. Figure 4.9 (b) shows an

asymmetric peak at 511.5 cm-1

with an FWHM of 15.8 cm-1

for sample Si5. The peak is red

shifted as compared to bulk silicon by 9 cm-1

and have an asymmetric nature, which can be

assigned to the phonon confinement effect in the NSs. Such shift has been predicted to be

arising from the confinement in a 2D planar structure of Si of thickness of the order of 2 nm

[15].


102

4.1.2.7 Investigations of SiNTs (Si5) with nano-beam electron energy loss spectroscopy

Figure 4.10 shows the experimental nanobeam low electron energy loss spectra

(NEELS) for two different NTs (curve (a) and (b)) and spherical NPs (curve (c)) compared

to the SiO2 one (curve (d)).

Figure 4.10 Nanobeam low electron energy loss spectra for two different nanotubes (curves (a) and

(b)), a spherical nanoparticle (curve (c)) and a SiO2 standard (curve (d)); inset: the complete

experimental SiO2 NEELS spectrum presenting the zero-loss peak due to elastically transmitted

electrons and first order plasmon features at energies between 10 and 30 eV.

The curves present a zero-loss peak due to elastically transmitted electrons (indicated

by the arrow in the inset of figure 4.10) and first order plasmon features at energies between

10 and 30 eV. The plasmon peaks of all, except one (a), are centered at about 14 and 23 eV,

which is the fingerprint of the SiO2 and shows the chemical composition of both NTs and

spherical NPs [16]. Instead, the curve (a) exhibits two features centered at about 12 and 23

eV, the former more intense than the latter. The peak at 12 eV is generally ascribed to the

so-called surface plasmon oscillations of clean Si [17]. No hints of feature at 17 eV, due to

the silicon bulk plasmon oscillation, is detectable: in fact a valley is present in the spectrum

(a) at this electron loss energy [17]. It suggests that this peculiar NT might be constituted by

a clean thin silicon layer and a very small amount of silicon dioxide. On the other hand, by

calculating the integrated intensity of the elastic peak, I0, and that of the first order plasmon

features, Ip, information on the nanostructure thickness, T, can be obtained through the

following relation [18]:


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ln (Ip/I0) = T/ 4.1

where, is the electron mean free path. Since maintaining the same experimental

conditions the value does not change, we calculated the ratio between the measured

diameter, d, and the T/ values of several NTs, spherical NPs and a SiNW. Then, by

assuming T/d of the SiNW equal to one, we obtained values ranging between 0.5 and 0.6 in

all the other cases, i.e. the nanostructure thickness T crossed by the impinging electrons is

well thinner than their morphologically measured diameter. In other words, this analysis not

only gets rid of any doubt on the nanotubular nature of the elongated structures but also

gives clear hints on the hollow nature of the spherical NPs.

Figure 4.11 Si L2,3 edge electron energy – loss spectra recorded for the SiO2 specimen (curve (a)), a

spherical nanoparticle (curve (b)), two nanotubes (curves (c) and (d)), and the clean Si nanotube

(curve (e))

Figure 4.11 shows a comparison among the typical Si L2,3 edge electron energy loss

spectra recorded for NTs (curves (c) and (d)), spherical NPs (curve (b)), the SiO2 specimen

(curve (a)) and the SiNT (curve (e)) reported in ref. [2] and [3]. The curves (a), (b) and (c)

show the silicon dioxide characteristic edge feature in the 105–110 eV region, suggesting

the presence of silicon dioxide as the chemical composition in the NPs as well as in the NT

(c). Such a peak is completely absent in spectrum (e) of the clean SiNT which exhibits the

typical feature (located between 100 and 104 eV) corresponding to a dipole transition from

the unoxidised Si 2p shell to the bottom of the conduction band [19]. The spectrum of some

NTs (curve (d)) exhibits a peak typical of SiO2 and a shoulder located at an energy loss

energy between 100 and 104 eV, which is the fingerprint of the clean silicon spectrum


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(curve (e)). A feature like this is completely absent in the Si L2,3 edge electron energy loss

spectrum for pure silicon dioxide (see the vertical dotted line in figure 4.11). This means that

the NT corresponding to curve (d) is not completely oxidized, although fully oxidized NTs

were more frequently observed than the others. These results were also confirmed by

evaluating the percentages of Si and O atoms from the analysis of Si L2,3 and O K edge

electron energy loss spectra [18], again indicating that, in the spherical NPs and in the major

part of the NTs, Si atoms are about 33% while O atoms reach 66% as for the SiO2 standard

sample. In the remaining cases, the percentage of Si atoms is, instead, a little bit higher,

around 44%.

4.1.2.8 Analysis of SiNTs (Si5) by HRTEM and nanobeam electron diffraction

Figure 4.12 EF-HRTEM image of a nanotube. The upper left inset reports the FFT of the area

contained in the white square; the lower right inset shows the filtered image of the region obtained

by making the inverse of the FFT displayed in the upper left inset.

In order to give a closer inspection to less oxidized NTs, we carried out high

resolution TEM images on some of the nanotubes showing the typical feature of the clean

silicon at the Si L2,3 edge spectrum. A typical EF-HRTEM image, acquired in the “Scherzer

defocus” conditions, is reported in figure 4.12. Very surprisingly, the nanotube appears as a

patchwork of locally ordered and highly disordered regions. In the upper left inset of figure

4.12, we report the texture analysis performed by the Fast Fourier Transform (FFT) of the

image of a particularly ordered region, showing a hexagonal pattern. Measuring the distance

between the spots and the center, a value of 2.60 ± 0.32 nm-1

is obtained, corresponding to

0.38 ± 0.05 nm in the direct lattice. In the lower right inset of figure 4.12, we show the

resulting region after having filtered it by applying a hexagonal mask to the six spots of the


105

FFT. This area clearly appears as an overlapping of two hexagonal atomic arrangements, to

be probably ascribed to the upper and lower lateral surface of the tube. Moving around the

imaged NT, spots in the FFT can rotate, double, increase in number and even elongate so to

have a resulting pattern not dissimilar from that of a high polycrystalline region. This

behavior reminds closely what has been observed by HRTEM on a graphene oxidized layer,

where highly atomically ordered and pure C areas alternate with disordered oxidized regions

[20]. Moreover, EF-HRTEM images of these NTs enable us to measure the thickness of

wall, which results to be around 0.5-0.7 nm. Interestingly, the thickness of an oxidized

graphene layer has been reported to range between 0.6 and 1.2 nm, mainly because of the

presence of oxygen atoms over and below the graphene sheet and, to a lesser extent, of the

distortion of the C network due to O absorption which induces a change of carbon atoms

hydridization from sp2 to almost sp

3 [21]. All these similarities suggest that we could be

facing with tubular structures formed by a mostly oxidized silicon layer. As a matter of fact,

it has been reported that a monolayer of amorphous silicon dioxide on a Si(111) surface can

be estimated to be as thick as 0.3 nm [22]. However, in the present case, both the internal

and the external surfaces of the tubes could be oxidized, so to increase the overall thickness

of the wall. On the other hand, we have to take into account that, evaluating the wall

thickness of a NT from a TEM image, we surely overestimate the thickness because of the

curvature of the NT and the image formation mechanisms [23]. For example, in the case of

CNTs, even if the C covalent radius in the sp2 hybridized layer is about 0.073 nm, the NT

wall thickness measured from the HRTEM image, in our apparatus, is about 0.2-0.3 nm. As

far as the FFT of the EF-HRTEM image of the highly ordered atomic regions is concerned,

the observed hexagonal pattern, shown in the upper inset of figure 4.12, paves the way for a

few considerations. Such a hexagonal spot arrangement can be originated by a quasi-two

dimensional hexagonal direct lattice, like that of graphene and CNTs, or to a (111) oriented

sheet (puckered layer) of a bulk diamond lattice, like Si bulk. In both cases, the observed

spots correspond to four {10} and two { } reflections. As a consequence, since the

distance d10 (or equivalently ) is equal to aHEX √3/2, where aHEX is the two-dimensional

hexagonal lattice constant, assuming d10 equal to 0.38±0.05 nm, our experimentally

measured value, would lead to an aHEX equal to 0.43±0.05 nm. The obtained value is

compatible with the (calculated or measured) lattice constant of both the sp2 hybridized

SiNT or silicene layer (0.38-0.43 nm) [24 – 35]. This means there is no chance to distinguish


106

between the two hybridizations. However, the unreactivity of these areas with oxygen can

only be explained by sp2 – hybridization [36] or Si-H bonds due to the presence of hydrogen

in the synthesis chamber.

From the analysis of the Si L2,3 edge it is worth noticing that the NTs are found to be

fully oxidized, and they do not exhibit any atomically ordered region at EF-HRTEM. As a

consequence, their FFT is similar to that of a highly polycrystalline specimen. NED results

support this structural model.

Figure 4.13 NED of the nanotube imaged in the upper left inset of the figure. The bright circular area

indicates the region from which diffraction pattern arises. In the upper right inset the profile of the

diffraction pattern obtained along a straight line passing through its center is reported.

Figure 4.13 reports a typical NED pattern recorded on a NT area as small as 50 nm

in diameter (Figure 4.13, upper left inset). The pattern is dominated by Debye (diffraction)

rings occurring due to the highly polycrystalline nature of the sample. The distance from the

center of the two main Debye rings is 2.36 ± 0.11 nm-1

and 7.6 ± 0.6 nm-1

, corresponding to

0.42 ± 0.02 nm and 0.13 ± 0.01 nm in the direct lattice, respectively. The same two Debye

rings and similar distance values within the errors have been measured for the amorphous

SiO2 standard.

The interpretation of all these results can be summarized as follows. In the presence

of a hydrogen and argon mixture, the synthesis process is able to form single wall SiNTs

characterized by a patchwork of two main types of regions. The former is an ordered two-

dimensional network of sp3 hybridized silicon atoms unable to react with oxygen, while the

latter, when exposed to air, prevalently oxidizes via modification of the original Si network.


107

The reasons why the atomically ordered areas do not suffer oxygen modification might

reside in the possible hydrogen bonding of Si atoms or the sp2- hybridized Si atoms

constituting these areas with respect to the other regions. The observation of atomically

ordered regions through EF-HRTEM in correspondence of less oxidized NTs is of

paramount importance because it suggests the existence of single wall SiNTs with sp3 or sp

2

hybridization which, up to now, has been only predicted by theoretical calculations.

Figure 4.14 (a) EF-HRTEM image of a nanoparticle; the inset shows the FFT calculated for the

white square region, and (b) NED of the same nanoparticle also imaged in the upper left inset of the

figure. The bright circular area indicates the region from which diffraction pattern arises. In the

upper right inset the profile of the diffraction pattern obtained along a straight line passing through

its center is reported.

Figure 4.14 (a) shows the EF-HRTEM image of a spherical NP from sample Si5. No

hint of highly ordered areas can be detected. This is also confirmed by making the FFT on

several small regions, scanned all over the NP. Nevertheless, sometime FFTs still present

some feature evoking a hexagonal pattern (Figure 4.14(a), inset). Besides, NED patterns are

very similar to those of NTs. All these results indicate that silicon NPs are highly

polycrystalline. On the other hand, the chemical composition analysis clearly points out

towards the presence of silicon dioxide nature of these nanostructures. Finally, their wall

thickness, directly measured from the EF-HRTEM images, amounts to a value between 0.5

and 0.7 nm, similar to that in the NTs. All these findings suggest that there are many

similarities with the highly disordered regions of NTs. Then, this means that we are dealing

with SiO2 hollow spherical NPs that could be the product of the oxidation of a single layer

Si spherical NP. Interestingly, such a single layer nanostructure is expected to form only in


108

case of sp2 hybridization. In this scenario our results are very peculiar: as a matter of fact, a

stable hollow nanosphere very similar to fullerene is not expected to form if silicon atoms

are sp3 hybridized. Nonetheless, to the best of our knowledge, no calculation has been

reported on the viability of such a silicon nanostructure. Reasons of the complete oxidation

of the hollow NPs with respect to partial one presented by some NTs could be related to

their spherical geometry unabling the formation the Si-H bonds.

4.1.3 Conclusions

From the analysis so far, it is clear that the NTs were highly oxidized, but the point

to be highlighted is that these are the evidence of formation of single layered structures in

silicon. Further, the experiments were carried out by increasing the content of hydrogen

during synthesis and this time the precursor with higher purity and treated for the removal of

the passivated layer of oxygen to avoid the content of oxygen during synthesis. The

observed results, however, did not show the formation of NTs, which indicated oxygen, also

played the important role in the formation of nanotubular structures. The results obtained by

increasing the hydrogen content are also discussed separately in detail. The exact role of

oxygen could not be discovered, but, possibly, the presence of oxygen decreases the melting

point of the precursor and oxygen might also be playing an important role in passivation of

single atomic layers similar to hydrogen, restricting their growth further. Finally, these

single atomic layers get curved in the form of nanotubes and nanoparticles to attain stability.

Another important point to be noticed relates to the stability of the structures formed. No

deformation in the microstructure was observed even after several months and even few

years as was inferred from the TEM observations.

4.2 Synthesis of silicon nanostructures in presence of different hydrogen

concentrations and its effect on the morphology


The SiNSs were synthesized following the method described in section 3.1.2. The

silicon powder, from Kemphasol (99 % Purity), India Ltd., was placed in the graphite

crucible that acted as anode. The details of the synthesis parameters are mentioned in Table

4.2. In this set of synthesis only the ambient gas ratio was varied keeping all the other


109

parameters of the synthesis constant. The arc current of 80 A, arc voltage of 12-16 V and

ambient pressure of 500 torr was maintained.

Table 4.2 The details of the synthesis parameters

Sample name Mole ratio (Ar/H2)

S1 (100/0)

S2 (95/5)

S3 (90/10)

S4 (85/15)

XRD patterns of the samples were recorded with Bruker D8 XRD machine with

CuKα radiations, Ni filter and graphite monochromator. TEM images were recorded by

Technai G2 twin TEM with a 200 keV LaB6 thermionic emitter and a CCD camera. Fourier

Transform Infrared Spectroscopy (FTIR) measurements were carried out with the resolution

of 2 cm-1

and averaging of 200 scans.


As referred earlier [7] enthalpy of the plasma increases with increasing hydrogen

content and is associated with a steep increase at 3000 K. This occurs due to the dissociation

of hydrogen, completing at 4000 K. The increase in the enthalpy of plasma leads to the

increase in the evaporation rate of silicon precursor and also changes the temperature

gradient of the plasma. The change in temperature controls the mechanism of nucleation and

the growth of SiNPs. Specific heat and thermal conductivity also increase with increasing H2

concentration [37], which is also accounted due to reasons discussed above. The increase in

thermal conductivity allows faster transfer of heat, hence again enhances the evaporation

rate, thus affects the size and type of the nuclei formed.


Figure 4.15 shows the XRD patterns for all the four as synthesized samples. The

XRD peaks match with the standard JCPDS Card no. 050565 corresponding to diamond Si.

It can be observed that each of the as synthesized samples consisted of XRD peaks

corresponding to Si(111), Si(220) and Si(311) planes. The XRD peak intensities is found to

be increased with increasing H2-concentration except for S4 in which there is formation of


110

β-SiC alongwith Si. The average crystallite sizes calculated from Scherrer formula is found

to be 8.5 nm, 13 nm, 16 nm and 14 nm for samples S1, S2, S3 and S4 respectively.

Figure 4.15 X – Ray diffraction pattern of samples synthesized in increasing H2 – concentration.


Figure 4.16 shows TEM micrographs of the samples synthesized in different gas

atmospheres. Sample S1 shows the presence of flake like nanoparticles of silicon and few

nanowires with diameters varying between 4 nm to 10 nm and the length of the order of 100

nm. Nanowires were found approximately to be less than 10 %. In S2, the number of

nanowires increased to nearly 60 %, flake like structures decreased and some large spherical

nanoparticles of size varying between 10 to 60 nm appeared. Sample S3 consisted of

spherical particles having tail of nanowires. The diameters of spherical particles varied from

10 to 100 nm while diameter of tail was found between 4 - 6 nm. Sample S4, which was

synthesized in 15 atomic % of H2, consisted of triangular and hexagonal platelets that

belonged to silicon carbide.


111

Figure 4.16 TEM micrographs of silicon nanostructures (a) S1, (b) S2, (c) S3, and (d) S4.

The observed results can be understood by looking at the enthalpies of the gas

compositions. When synthesis was done in presence of argon alone, because of the low

enthalpy of plasma the evaporation rate of silicon was low. This resulted in growth of

undefined nanostructures with poor crystallinity, which soon got oxidized on exposure to the

atmosphere. Later with an increase in hydrogen more crystalline forms were obtained and

resulted into the growth of one dimensional structure. This growth is possibly due to

formation of small nuclei that get condensed. Similar nuclei approach each other at the time

of condensation to form one dimensional structure. This is also visible from HRTEM images

of silicon nanowires (Figure 4.17) that show twin boundaries after a certain length.


112

Figure 4.17 TEM images of silicon nanowires; (a), (b), (c) and (d) magnified images of the squares

shown by arrows; (a) and (b) show the lattice spacing on silicon nanowires showing twin boundary,

(c) the lattice spacing on silicon nanowires, and (d) lattice spacing at the mouth of a nanowire.

Figure 4.17 (a) and (b) clearly shows the twin boundary (marked with black coloured

rectangles) where the orientation of the planes is changing. The lattice spacing of ~3.25 Å

corresponding to Si (111) plane (lattice spacing for bulk Si (111) is 3.11 Å) is observed. The

increased lattice spacing value may be attributed to the lattice dilation resulting from

reduced dimension. Figure 4.17 (b) shows the HRTEM image of a nanowire that shows the

lattice spacing of ~3.25 perpendicular to the axis of a wire near the tip and the planes grow

at an angle of 116° to these after a length of 15 nm (boundary marked by a black coloured

rectangle in figure 4.17 (b)). Similar lattice can be observed in figure 4.17 (c). Lattice

spacing observed at the tip of the nanowires is ~1.94 Å that corresponds to the interplanar

distance between Si (220) planes.

Figure 4.18 shows the high resolution images of spherical structures observed in

sample S3. The lattice observed on the surface of a sphere is 3.18 Å (Figure 4.21(b)) and

corresponds to the interplanar distance between Si (111) planes. This type of structure

formation might be again due to increased enthalpy, which allows high evaporation rate and

the growth of nanoparticles. The lattice spacing here has come close to bulk silicon as the

particles are large. The spherical particles thus formed are single crystalline in nature.


113

Figure 4.18 (a) TEM image of spherical nanoparticles of silicon observed in Sample S3, (b)

magnified image showing lattice planes and (c) fast Fourier transform of image (b).

Figure 4.19 HRTEM image of hexagonal platelet of silicon carbide; lower left inset shows the

magnified image of the region marked by square and lower right inset show the corresponding fast

Fourrier transform

Figure 4.19 shows HRTEM image of one of the hexagonal platelet found in the

sample S4 that bears lattice spacing of ~ 2.57Å. This lattice spacing corresponds to (101)

plane of hexagonal silicon carbide. The planar dimensions of the sheet varied from 20 nm to

150 nm. Although the thickness could not be exactly calculated, it can be said that the third

dimension is much smaller than the other two. The high enthalpy due to the presence of the

15 % H2 has made carbon to evaporate from the crucible and react with the silicon vapours

to form SiC.


114

4.2.2.3 FTIR analysis

FTIR spectra were recorded to confirm the species present in the samples. The

results are shown in figure 4.20.

Figure 4.20 FTIR spectra of silicon nanostructures synthesized in different gas compositions.

Absorption due to Si-O-Si stretching, bending and rocking modes are observed

around 1090 cm−1

, 812 cm−1

and 463 cm−1

respectively [38,39] while an absorption band

around 2250 cm-1

due to the Si-H stretching mode is also seen. As the concentration of

hydrogen gas during synthesis increases, the relative absorption of the bands for Si-O-Si

bond decreases and that for Si-H increases. When hydrogen concentration is increased to

15% the absorption for Si-H band has decreased while a sharp peak for SiC appears at 800

cm-1

. The broad band observed around 3350 cm-1

is due to Si-O-H absorption, which is most

intense in sample S2. The oxygen observed in the samples might be due to exposure of

samples to atmosphere and some amount of oxygen and moisture may be expected in the

chamber as the evacuation is done up to 10-3

mbar. The sample synthesized in presence of

argon is not hydrogen capped and so is more prone to oxidation. In S2, Si-O-H absorption is

greater due to the presence of hydrogen during synthesis. Even the surface seems to be

capped with hydrogen. In case of S3, Si-H absorption is maximum and Si-O-H stretching

peak intensity has reduced. In S4, Si-H bond has disappeared while Si-C is observed at 800

cm-1

. This might be due to evaporation of carbon taking place from the crucible that reacts

with hydrogen forming gaseous hydrocarbons and also with Si to form SiC. Further work on

SiC was inspired from the formation of SiC here.


115

4.2.3 Conclusions

The desired morphologies of SiNSs can be obtained by changing the amount of

hydrogen during synthesis. The formation of different morphologies was justified based on

the thermodynamic behavior of plasma due to changing H2 concentration. Technologically

important SiC can be obtained after increasing H2 content beyond a certain concentration.

4.3 Antibacterial study of silicon nanoparticles (Si1) and nanotubes (Si5)

4.3.1 Introduction

Nanomaterials, due to their large effective surface area with potent number of

reactive sites, are extensively used for biological applications which include antimicrobial

activities, drug delivery and medical imaging [40 – 42]. Study of antimicrobial activity of

the nanomaterials is essential for biological applications especially to develop antimicrobial

medicines [43] and to find the suitability of a material for antimicrobial surfaces [44 – 47].

This is important since increasing resistance of microorganisms to multiple antibiotics has

raised the demand of effective, resistance free, cheaper and biocompatible antimicrobial

agents.

Nanomaterials provide a novel way to replace antibiotics. NPs, for instance have

been used to reduce skin diseases [48,49] and to prevent the microbial colonization formed

on the surface of devices like endotrachial tubes, catheters and prostheses [50,51]. Although

NPs of metal (silver, gold etc.) and some oxide systems (like TiO2, ZnO, etc.) exhibit

convincing antimicrobial activity, they have poor dispersion stability in the organic medium

[52]. This pose difficulties for further application like fabricating antimicrobial polymer

composites. The drawback with silver NPs lies in its rapid oxidation and ease of

agglomeration that changes its reactivity to the environment [53]. Besides, it is an expensive

material and its environmental hazards are a matter of concern [54].

Nanostructures of silica, on the other hand, are preferable over other nanomaterials

due to its noncorrosive and biocompatible properties. Also, they are chemically stable and

are not as expensive as silver and gold. The surface of silica structures can be functionalized

by many ways that make its integration easier with polymers. Silicon derivatives with

polymers have been used as anti corrosive and chemical resistant coatings [51].

Nanostructures of silica can thus impart much useful antimicrobial properties to the


116

polymers [55]. Song et al. [56] have reported the antimicrobial studies of silica NPs

modified by polymer whereas Hebalkar et al. [57] have reported the synthesis of silica NPs

for antibacterial and self-cleaning surfaces [57]. NPs of Silica are used in biological

applications [58,59] as scaffolds for drug delivery by functionalizing their surfaces and

attaching bio molecules [60,61]. Likewise, NPs of silica are used to form composites with

metal NPs like silver to avoid their aggregation and hence enhance antibacterial activity

[62]. They are also important with a perspective of dental applications, mainly as fillers [63].

Further, nanosilica exhibiting hollow structure is supposed to be superior compared

to the solid structures as hollow nanostructures possess greater effective surface area. They

are expected to provide increased effective surface reactivity required for charge

accumulation that can contribute to antimicrobial activity. Here, the as synthesized NTs

were mostly oxidized thus very nearer to silica. So, the antimicrobial activity of these NTs

was investigated. The activity of NTs (sample Si5) was compared with the activity of

sample Si1, consisting of SiNWs and SiNPs, with surface oxidized.

The antimicrobial activity of as synthesized nanostructures was investigated for the

selected strains of bacteria which are pathogenic and frequently colonize the medical

devices. The Gram-positive bacteria included were Bacillus subtilis and Staphylococcus

aureus whereas Gram-negative bacteria were Escherichia coli and Pseudomonas

aeruginosa. S. aureus is one of the major resistant pathogens found on the mucous

membranes and the human skin of around one-third of the population and it is extremely

adaptable to antibiotics [64]. The antibacterial activity was assayed using optical

densitometry technique and viable cell counting method using plating technique.


4.3.2.1 Bacterial culture and growth

The bacterial strains were procured from National Collection of Industrial

Microorganisms (NCIM), India. The Gram-positive bacterial cultures tested were

Staphylococcus aureus (NCIM 2079) and Bacillus subtilis (NCIM 2063). The Gram-

negative cultures used were Escherichia coli (NCIM 2065) and Pseudomonas aeruginosa

(NCIM 2200). These bacteria were cultured in Nutrient Broth (NB) media. (For the

preparation of NB media, 10 g Yeast extract, 5 g sodium chloride, 10 g tryptone, is

dissolved in 900 of millipore water in a conical flask. The pH of the medium is maintained


117

between 7-7.5. This pH is essential for optimum growth of the bacteria used in this study. 20

g Agar is then added to the above mixture and the solution is autoclaved (121°C/15psi).)

Loopful of cultures were pre inoculated in 2 ml of NB media. These cultures were then

grown at 30oC overnight (B. subtilis) and at 37

oC overnight (all other bacteria). Except B.

subtilis, the population of microbes was maintained between 1 x 108

and 5 x 109

CFU/ml

(CFU means Colony Forming Units).

4.3.2.2 Antibacterial test

I. Estimation of minimum inhibitory concentration (MIC) using optical

densitometric technique

The micro dilution method was utilized for estimation of the MIC of the samples for

determining the antibacterial activity. The colloidal solutions of samples Si1 and Si5, with

concentration of 1mg/ml were prepared by ultrasonicating them with distilled water for 30

min. The experimental miniprep was prepared having following components: 900µl of NB

media, 100µl of respective bacterial inoculums, samples Si1and Si5 in varying

concentrations (0, 10, 50, 100, 150 & 200 µg/ml) and sterile distilled water was added to

equate the reaction volume. The experimental miniprep were allowed to grow overnight at

30oC (for B. subtilis) and rest of the strains were incubated at 37

oC over night. Each and

every test concentrations along with control was diluted up to 109

dilutions using 0.9%

sodium chloride solution.

To determine the MIC values, 100 μl of all test concentrations and their dilutions

were added to 96 well plates and absorbance at 600nm was recorded. NB media and Saline

were used as negative control. Experimental mixture without SiNSs (100 μl) was used as

positive control. Both sample Si1 and Si5 at different concentration without bacterial

inoculums and their dilution upto 109 in saline was tested for its absorbance interference.

(Optical densitometric technique is based on the idea that light passing through a

suspension of microorganisms is scattered, and the amount of scatter is an indication of the

biomass present in the suspension. If the concentration of scattering particles becomes high,

then multiple scattering events become possible. Light scattering techniques to monitor the

concentration of pure cultures have the enormous advantages of being rapid and

nondestructive. However, they do not measure cell numbers nor do they measure CFU.

Light scattering is most closely related to the dry weight of the cells.)

Chapter 4. Synthesis of silicon nanostructures and their applications

*Due acknowledgement to Dr. Sujatha Raman, Prof. S. Gosavi and Prof. W. N. Gade 118

II. Determination of CFU

In order to determine the CFU spread plate method was adopted. In this method,

nutrient agar plates are prepared by pouring hot agar solution in the petridishes and allowing

them to cool. The nutrient agar provides essential nutrients for the growth of bacteria. The

bacterial suspension is then poured with the help of a micropipette at the centre of the agar

plate. A glass spreader is used to spread the poured inoculums uniformly on the nutrient agar

plate. The glass spreader is heated in the flame of a Bunsen burner and then dipped in

ethanol to remove the unwanted bacteria present on the spreader. These plates are then kept

at 37°C for 24 hours and the bacteria are allowed to grow.

In our case, 100 µl of the experimental miniprep was used from each and every test

concentrations having dilutions of 109, 10

8 & 10

7 to plate the bacterial inoculums including

control. In B. subtilis, plating was carried out in dilutions of 103, 10

4 & 10

5 owing to low

optical densitometric measurements. The colony forming units were determined by counting

the bacterial colonies and then by multiplying with the dilution factor. All assays were

carried out in duplicate. These experiments were carried out at Department of

Biotechnology, SP Pune University*.


4.3.3.1 Optical densitometric analysis

The values of Minimum Inhibitory Concentration (MIC) for samples Si1 and Si5

were obtained from the measurements of optical density at 600 nm for the inoculums

containing different bacteria. These are shown in figure 4.21 (a) and (b) respectively. There

are certain limitations associated with using optical density techniques to determine bacteria

viability in the presence of nanomaterials since they themselves contribute to optical density

at different concentrations [65]. To resolve this issue, optical densities of the nanostructured

samples Si1 and Si5 are provided along with the data for reference.


119

Figure 4.21 Effect of different concentrations of silicon nanostructures on bacterial strains tested.

Standard absorbance values of silicon nanostructures at various concentrations from 0 µg/ml

(positive control) to 200 µg/ml are provided. Experimental mixture having NB media with respective

bacterial inoculums, without nanostructures was used as positive control. NB media alone was used

as negative control (a) the effect of Si1 on bacterial strains tested, (b) the effect of Si5 on bacterial

strains tested.

It was found that the concentration of NPs at which the growth was inhibited was

different for different bacteria. In Si1 treatment, the growth of E.coli and B. subtilis was

inhibited at 10 µg/ml, while for P. aeruginosa, inhibition was observed at 50 µg/ml. S.

aureus did not show any inhibition by use of Si1. In contrast, Si5 showed MIC of 10 µg/ml

for S. aureus. IC-50 (inhibition of bacteria to 50 % of the untreated value) is found to be 100

µg/ml. The reports from “The Center for Disease Control and Prevention” indicate that the

number of annual Multidrug-Resistant Staphylococcus aureus (MRSA) infections increased

from 1,27,000 to 2,78,000 between 1999 and 2005 [66]. In this scenario, NT sample is found

to be a potential candidate to target MRSA infections. Inhibition of E.coli and P.

aeruginosa by Si5 is comparable with Si1. The growth of B. subtilis cultures were not

inhibited by Si5 in contrast to Si1. Thus, the data depicts the difference in the mechanism of

inhibition by these two nanostructures towards bacterial strains.

The inhibition occurs generally via reactive oxygen species inhibition, membrane

disruption, protein inactivation, flocculation or unknown mechanisms [65]. Further studies


120

are needed to be carried out to understand the mechanism of inhibition by these

nanostructures (Si1 and Si5). The MIC was found to be of the order of microgram for these

tested nanomaterials, which is comparable to those reported in metal oxides for both Gram-

positive as well as Gram-negative bacteria [65]. In many cases, MIC values obtained for

these nanostructures are better than reported in other oxide systems [65].

4.3.3.2 Colony forming units analysis

In order to measure the viable cells, colony forming units were determined using

serial dilutions of suspensions followed by spread plate colony counting. Figure 4.22 (a) and

(b) shows the plot of CFU obtained for B. subtilis for samples Si1 and Si5 respectively while

(c) and (d) for S. aureus for samples Si1 and Si5 respectively.

Figure 4.22 Colony forming units counting in Gram-positive bacterial strains calculated for different

concentrations of nanostructures (0 to 200 µg/ml) (a) CFU of B. subtilis cultures calculated at the


4 and 10

5 for Si1, (b) CFU of B. subtilis cultures calculated at the dilutions of 10

3,

104 and 10

5 for Si5, (c) CFU of S. aureus cultures calculated at the dilutions of 10

7, 10

8 and 10

9 for

Si1, and (d) CFU of S. aureus cultures calculated at the dilutions of 107, 10

8 and 10

9 for Si5.

When exposed to different concentrations of Si1, B.subtilis showed reduced viability

at 100 µg/ml (Figure 4.22 (a)) whereas increase in the viability is observed at 200 µg/ml.

This may be attributed to the interaction of sample Si1 with bacteria at high concentration.

The B. subtilis cultures, exposed to Si5, showed a definite pattern of reduced viability

(Figure 4.22 (b)) with increase in sample Si5 concentration; although initial optical

densitometry analysis could not clearly delineate inhibition (Figure 4.21(b)). The IC-50

value of Si5 was 200 µg/ml in B. subtilis cultures.


121

In S. aureus (Figure 4.22 (c)), the IC-50 value was found to be 100 µg/ml for Si1.

However, the optical densitometric analysis could not reveal the inhibition (Figure 4.21 (a)),

possibly due to the interference of the optical activity of sample Si1. The 10 µg/ml of Si5

proved to be effective in controlling the S. aureus even at very low concentration with IC-50

of 100 µg/ml (Figure 4.22 (d)). Although nanomaterials like ZnO [66], Fe3O4 [67] and Ag

particles were proved effective against MRSA infection, NTs sample Si5 were proved to be

effective even at very low concentrations (10 µg/ml). Thus, the biocompatible nature and the

cost-effective, eco-friendly synthesis adds to the credit of oxidized silicon NTs as effective

antibacterial agent.

Figure 4.23 (a) and (b) shows the plot of CFU obtained for E. coli for samples Si1

and Si5 respectively. Figure 4.23 (c) and (d) shows the plot of CFU obtained for P.

aeruginosa for samples Si1 and Si5 respectively. With Gram-negative bacteria like E. coli,

MIC was found to be 10 µg/ml for both Si1 and Si5, which is in concurrence with

densitometric analysis (Figure 4.17 (a) and (b)). A definite pattern of inhibition was

observed in both Si1 and Si5 with the increase in concentration. In P. aeruginosa, 50 µg/ml

was found to be effective in reducing the viability both in Si1 and Si5 (Fig. 7 c and d). Thus,

Si1 and Si5 are found to be competent in controlling both Gram-positive and Gram-negative

bacterial strains tested.

Figure 4.23 Colony forming units counting in Gram-negative bacterial strains calculated at the


8 and 10

9 for different concentrations of nanostructures (0 to 200 µg/ml) (a) CFUs

of E-coli cultures for Si1, (b) CFU of E-coli cultures for Si5, (c) CFU of P. aeruginosa cultures for

Si1, and (d) CFU of P. aeruginosa cultures for Si5.


122

4.3.4 Conclusions

It is shown that oxide coated silicon nanostructures can be a good substitute as a

cheaper and biocompatible antimicrobial agent. The low values of MIC are encouraging and

they point out towards the specific surface properties of the silicon nanostructures. It is

proposed that on account of the extremely thin oxide layers of the silicon nanostructures, the

interaction of these with the bacteria becomes strong and capable of inhibition. The role of

NTs in controlling the MRSA infections is emphasized in this study, making it efficient

antiMRSA agent.

4.4 Field emission study of silicon nanotubes (Si5)

4.4.1 Introduction

Field emitters are used in flat panel display technology as well as in the cold cathode

technology for electron tube devices such as microwave tubes. The conventional thermionic

emission devices are accompanied with high power dissipation due to high cathode

temperature. Nanomaterials are important as the nanodimensions make them suitable

candidate for field emission. Several metallic and semiconducting nanomaterials are found

to operate at a lower applied potential delivering high current density as compared to

conventional field emitter counterparts. The nanowire forms of various materials are suitable

for all observed field emission properties. The study of field emission from SiNSs is

important as they can be merged with the existing Si technology. Literature survey shows

that field emission (FE) study on various SiNSs, synthesized by different routes like electron

beam annealing [68], thermal evaporation using vapor-liquid-solid (VLS) mechanism [69],

laser ablation technique [70] etc., have been carried out. Tubular structures are again more

important because of reduced wall dimensions. But, SiNTs are not much explored for field

emission owing to the difficulty in its synthesis. Time-dependent density functional study

reports carbon-like silicon nanotubes, as better field emitter than its carbon and boron nitride

counterparts [71]. So, exploring silicon nanotubes for their field emission (FE) studies and

other applications was important.

4.4.2 Electron field emission

Electron field emission (FE) is emission of electron from a condensed phase to

vacuum under the influence of strong electric field. FE is a direct realization of quantum


123

mechanical phenomenon of electron tunneling when the potential barrier becomes

comparable with the de Broglie wavelength of the electron. The emission of electron from

cold metal upon application of high electric field was first observed by Schottky in 1923.

The theoretical explanation was proposed by Fowler and Nordheim in 1928. Fowler and

Nordheim proposed the F-N equation which explained the field emission behavior for a

single emitter. But, as in present case, a modified Fowler Nordheim equation [72]

is adopted

for multiple emitters. The modified FN equation is given as Eq.(1)

,

(5.2)

where, J is the current density E is the local/surface electric field, a and b are constants (a =

1.54 x 10-6

AeV V-2

, b = 6.83 eV-3/2

Vnm-1

), is the work funtion of the emitter and is the

field enhancement factor. The electron emission is a function of suraface states which

depend on the local electric field (Eloc) which in turn depends on the geometry of the emitter.

The factor connecting the applied electric field (Eapp) and the local electric field is termed as

field enhancement factor, .

4.4.3 Experimental procedure for field emission study

To study field emission (FE) properties of silicon nanotubes sample the sample was

treated with HF to remove the surface oxide layer. The powder was then pasted on a

tungsten blunt tip, having diameter ~ 100 µm using conducting silver paste. The blunt tip

was prepared by itching tungsten wire (diameter 0.3 mm) in KOH solution. The tip was then

mounted on a copper rod attached with linear motion drive which facilitates change inter

electrode distance. The copper rod is mounted in all metal chamber for FE study. The

emitter serves as cathode while phosphor coated ITO glass acts as anode, forming a planer

diode configuration. The chamber is equipped with rotary backed turbo molecular pump to

attain a pressure upto 10-5

– 10-6

torr. FE studies require ultra high vacuum (UHV). Hence,

to further improve the vacuum, the chamber is baked for 8 hours at 200°C. The chamber is

then pumped using sputter ion pump and titanium sublimation pump. Finally, a pressure of ~

10-8

torr is attained. Now the potential difference between the two electrodes is increased.

The field for which electron emission starts (observed as fluorescence on phosphor coated


*Due acknowledgement to Padmashree Joshi and Prof. D. S. Joag 124

ITO glass) is called turn on field. In our case we have defined turn on field as field required

for attaining a current density of 10 µA/cm2. The emission current stability was monitored

using computer controlled data acquisition system with a sampling interval of 10 seconds.

The field emission micrographs are recorded using a digital camera (Canon SX150IS).

These experiments were performed in Field Emission Lab, Department of Physics, SP Pune

University.*


Figure 4.24 (a) shows the SEM image of the SiNTs coated W-tip. Figure 4.24 (b))

shows a plot of current density (J) verses applied electric field (E). The FN plot, which is a

graph of natural logarithm of (J/E2) verses (1/E), shown in the inset of figure 4.24 (b), is

almost linear in nature. A maximum current density of 4.2 mA/cm2 is attainable at an

applied electric field of 2.8 V/µm (inter electrode separation of 3000µm). The turn on field

defined to draw a current density of 10 µA/cm2 is merely 1.9 V/µm. The emission current

density measured for different applied electric fields is the total current emitted by the

randomly oriented nanotubes.

Figure 4.24 (a) SEM image of sample Si5 coated W- tip, (b) J-E plot (inset shows FN plot), (c)

emission current vs. time plot, and (d) FEM micrograph.

Figure 4.24 (c) shows the plot of emission current stability and figure 4.24 (d) the FE

image. It is observed that after certain period of time (~1 hour) the current lowers from its

preset value to ~ 0.7µA and then stabilizes. The ions/gas molecules are known to get


125

adsorbed or desorbed at the emitter surface. In particular, the puckered surface of the SiNTs

can provide many adsorption sites. The ionization of the residual gas molecules in the UHV

chamber can occur due to the applied electric field between the two electrodes. The ion

bombardment may extract adsorbed atoms at the emitter surface. Both these phenomena

occur simultaneously at the emitter surface leading to creation or destruction of certain

emitting sites. In due course of time, if the probability of formation of new sites is more than

that of destruction, then, the current is observed to increase than its preset value. However,

in the present study, certain emitting sites have been possibly destroyed leading to a

reduction in current. Thus, the adsorption-desorption phenomenon occurring at the emitter

surface i.e. the ambient atmosphere governs the observed FE current fluctuations. Ignoring

the initial fall in the preset current, the standard deviation was calculated to be ~ 9.7%.

Table 4.3 Comparative field emission study on Si nanostructures

A comparison with existing literature is interesting (Table 4.3). The β factor in the

present case has been estimated to be 5534, which is realistic. If the turn on field is

considered, the SiNTs synthesized by arc plasma method possess the lowest value except for

the SiNWs deposited by CVD [69,79]. The current density of 4.2mA/cm2 at 2.8 V/µm also

represents a good figure of merit. More importantly the FE current stability from SiNTs is

found to be good.

Morphology Turn on field β Jmax

SiNWs on C-cloth [69] 1.1V/µm (J =1mA/cm2) ~2-6x10

4

SiNWs on Si wafer [73] 3.4 V/µm (J = 1mA/cm2)

Self assembled SiNSs [68] 2.5V/µm (I = 1nA)

Two tier SiNSs [74] 10-14V/µm (J = 0.01

mA/cm2)

Vertically aligned SiNWs [75] 0.8MV/m (J = 10µA/cm2) 455 442µA/cm

2

(14V/µm)

Si nano-micro wire on Si

substrate [76]

15V (I = 10nA) 1µA (1150V)

Si thin film [77] Eth= 2.5V/µm (I = 1nA) ~18000

Self assembled SiNSs on n- &

p-type Si [78]

Eth= 2V/µm (I = 1nA)

SiNTs [present study] 1.9V/µm (J=10µA/cm2) 5534 4.2mA/cm

2

(2.8V/µm)


126

Here, it may be noteworthy to comment about the thin walls of the tubes as well as

the surface structures of the emitting nanotubes. The thin corrugated surfaces, resulting from

the puckered atomic arrangement of Si, as shown in the schematic diagram in figure 2.24

(b), might also be assisting in the FE properties leading to the moderately high value of the β

factor in spite of the crisscross arrangement of the tubes. Moreover, the part of the tube

surfaces with high surface states resulting from the dangling bonds or hydrogen bonds may

be the cause for the adsorption/desorption processes

4.4.5 Conclusions

SiNTs were subjected to FE studies at the base pressure of ~10-8

mbar. A maximum

current density of 4.2 mA/cm2 is attainable at applied electric field of 2.8 V/µm. A low turn

on field of merely 1.9 V/µm is required to draw a current density of 10 µA/cm2. The current

stability at 1 µA preset value is found to be good. The SiNTs are, thus, a potential candidate

for future application as a FE source [76].

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Chapter 5

Synthesis of Silicon carbide

Nanostructures & Application

This chapter provides the experimental outcomes concerning the efforts carried out in the

synthesis of SiC nanoparticles. Furthermore, some preliminary results regarding the composites

of SiC nanoparticles with DGEBA Epoxy are discussed in brief.

Chapter 5.Synthesis of silicon carbide nanostructures & application

131

5.1 Introduction

Silicon carbide is an interesting material on virtue of its properties that has created

applicability of its nanostructures (NSs) in various fields. The thermal plasma is one of the

well-known routes used for its synthesis, which yields the product in one step. In this

process different types of sources (solid or gaseous) can be used. Some researchers have

used gaseous precursors like silane [1], methane [1] and silicon tetrachloride [2] for the

synthesis. There are reports on the synthesis of SiC nanoparticles (SiCNPs) by using micron

sized particles of SiC [3], but methane had to be used for avoiding Si impurities [3].

Gaseous precursors, on the other hand, are difficult to handle, hazardous to environment and

not cost effective. Hence, some authors have used solid precursor combinations of Si [4–6]

or SiO2 [6,7] with C [4,5,7]. Even micron sized SiC have been used. However, the size

distribution of SiCNPs was wide and often possessed the impurities of both Si and C. Nayak

et al. [8,9] have used rice husk for synthesizing SiC, but the method has the disadvantage

that it leads to the presence of impurities [8].

Here, we aim at synthesizing SiCNPs, by arc plasma assisted synthesis using

microcrystalline particles of silicon and graphite as precursor. A controlled heating was

maintained by selecting the geometry of crucibles such that the heat was optimum to

evaporate Si and C together and react with each other to form SiC. However, for a certain

optimum conditions we could fully eliminate Si impurities during synthesis and C impurities

could be later removed simply by calcination in air without exercise of any chemicals. The

synthesis and study of SiCNPs forms the first part of this chapter.

Second part of the chapter consists of the study of the composites of Diglycidyl

Ether Bisphenol A (DGEBA) epoxy with SiCNPs. DGEBA is widely used as adhesive,

coating and encapsulated materials, due to its good mechanical properties and attractive

chemical and electronic properties [10]. As such, it finds a wide range of applications in

products like paints, surface coatings, adhesives, and electrical accessories. However, it

seem to suffer from major drawbacks in terms of poor resistance to crack initiation and low

impact strength. The great majority of the studies involve the chemical modification of

epoxy resin like with reactive liquid rubber [11,12], allyl glycidyl ether [10] and 2,3-

epoxypropyl methacrylate [10]. The other approach is by use of inorganic fillers like glass


132

[13], carbon [14], asbestos, oxides and textile fibers for improving the tribological

performance. The reduction in wear rate is mainly due to preferential load support of the

reinforcement components, by which the contribution of abrasive mechanism to the wear of

materials is highly suppressed. The micron sized particles bear disadvantages like

requirement of large amount of fillers, disintegration of the fillers and detached particles

[15] and high concentration of fillers is also detrimental to the processibility of polymers.

Thus, the use of nano fillers is an optimum alternative and they have proved to be better

[16–18].

SiC bears all the advantageous properties as mentioned above and hence it is used as

fillers in many polymers to improve its thermal and tribological properties [19–21]. Rodgers

et al. [22] have observed improved thermal and mechanical behavior in modified DGEBA

due to nano-SiC incorporation. Ji et al. [23] have used surface modified SiCNPs and found

better performance than non modified NPs. Similarly, Luo et al. [24] have observed

improvement in tribological behavior. Zhou et al. [25] have studied the thermal conductivity

of epoxy resin by the addition of a mixture of graphite nanoplatelets and silicon carbide

microparticles. Here, our aim was to improve the thermal and crack resistance properties of

DGEBA. Thus, SiCNPs were employed in the present work to prepare wear resisting

nanocomposites. But, SiCNPs could not be dispersed into DGEBA directly, so a method of

dispersion was found out using intermediate solvent and SiCNP-DGEBA composites were

prepared as a preliminary work and their properties were studied.

5.2 Synthesis and characterization of SiC nanoparticles


The graphite electrodes as described in section 3.1.1.2 were used for the synthesis of

SiCNSs. The source of silicon consisted of 99% microcrystalline (300 mesh) powder of

silicon from Sigma Aldrich while source of carbon consisted of microcrystalline graphite

powder from Kemphasol. The powders of the two materials were mixed in different

proportions and these served as the precursors. The arc voltage was maintained in the range

of 12-14 V and the cathode diameter was kept 0.9 cm at an ambient pressure of 500 torr

during the synthesis of all the samples. The ratio of silicon to carbon source used in

precursor during synthesis, which was changed for different samples in order to optimize the


133

conditions to achieve the desired results, is mentioned in Table 5.1. The synthesis was

carried out at two different arc currents of 80A and 100A for each crucible shape. Later, for

crucible CR5, the ratio of Si:C precursor was varied at 80 A with a view to improve yield

and reduce impurity. Thus, analysis presented will bear comparison in results for samples

synthesized (i) at same arc currents as a function of different crucible shapes, and (ii) for

each crucible shape as a function of arc current. The nomenclatures of the samples are also

listed in the table 5.1 for each combination of the experimental parameters of the samples

prepared. The synthesis was undertaken by following the procedure described in section

3.1.2.

Table 5.1 Details of the synthesis parameters used for the synthesis of SiC- nanoparticles.

Anode Diameter

(Crucible)

Sample

No.

Arc

Current

Precursor Ratio

(Atomic) (Si:C)

Gas

3.0 cm (CR2)

(Conical cavity)

SiC1 80 A 1:0

Ar:H2

(95:5) SiC2 100 A

3.0 cm (CR3) (Two

stage cylindrical

cavity)

SiC3 80 A

1:1

Ar

SiC4 100 A

1.7 cm (CR4) (Two

stage cylindrical

cavity)

SiC5 80 A

SiC6 100 A

1.0 cm (CR5)

(Two stage

cylindrical cavity)

SiC7 80 A

SiC8 100 A

SiC9 80 A 3:2

SiC10 80 A 7:3

Our main focus was (anode) crucible geometry which changed the nature of cooling

and helped in selective evaporation of the precursors. While performing experiments, it was

observed that the arc became more stable when the diameter of anode was reduced and it

approached towards the diameter of the cathode. The data generated from the experiments

showed interesting results about the amount of Si and C-impurities that varied with crucible

shapes and arc current. This was studied with the help of X-Ray Diffraction (XRD) and

Thermo gravimetric analysis (TGA). The morphology of nanoparticles was observed by

Transmission Electron Microscope (TEM) and lattice spacing by High Resolution TEM

(HRTEM). XRD patterns of the samples were recorded with Bruker D8 XRD machine with


134

CuKα radiations, Ni filter and graphite monochromator. TGA of the samples was carried out

from 30°C to 1000°C at the heating rate of 10°C/min rise, by passing oxygen at the rate of

80 ml/min using (Metler Toledo TGA 1). TEM images were recorded by Technai G2

ultratwin TEM with a 200 keV LaB6 thermionic emitter and a Charged Couple Device

(CCD) camera. For recording TEM images, the samples were first dispersed in isopropyl

alchohol by sonicating in an ultrasonic bath. Two to four drops of these dispersions were

then poured on the holy carbon coated copper grid (mesh size 200).


5.2.2.1 Yield of product

After the synthesis process, the NPs are collected from the upper part of the

chamber. When the precursor (anode) is evaporated and the NPs are formed due to super-

cooling, the lighter particles get deposited on the upper part of the chamber, while the larger

particles have a tendency to settle in the lower parts of the chamber. Some part of the anode

is also lost in the form of sputtered chunks of precursor material. Change of weight of

cathode takes place due to evaporation of carbon atoms from cathode and in some cases due

to deposition from the evaporated precursor of anode. For different crucible shapes, size and

arc current, rate of change of weight of anode (Awl) and cathode (Cwl), yield (‘Y’ is the

amount of product deposited on the upper part of chamber per minute) and ratio of Y to Awl

(expressed as β) obtained are primary observations that can be correlated with the results

obtained from other characterization techniques. These are expressed in the form of

formulae given by,

(5.1)

(5.2)

(5.3)

(5.4)

These observations were recorded and are presented in the form of graphs. Figure 5.1

(a) and (b) show the plot of Awl and Cwl for plasma current of 80 A and 100 A respectively

while figure 5.1 (c) and (d) show the plot of Y and β for plasma current for 80 A and 100 A

respectively.


135

Figure 5.1 (a) Plot of rate of change in weight of anode (Awl) and cathode (Cwl) for samples

synthesized using different crucible shapes at 80 A arc current, (b) plot of Awl and Cwl for samples

synthesized using different crucible shapes at 100 A arc current, (c) plot of yield (Y) and ratio of Y to

Awl (β) for samples synthesized using different crucible shapes at 80 A arc current, and (d) plot of Y

and β for samples synthesized using different crucible shapes at 100 A arc current.

Figure 5.1(a) shows that for the arc current of 80 A, the Awl was 22 mg/min for

SiC1 (crucible CR2, precursor Si), increased for SiC3 (crucible CR3, precursor Si:C=1:1)

and SiC (crucible CR4, precursor Si:C = 1:1) and again decreased for SiC7 (crucible CR5,

precursor Si:C = 1:1). For crucible CR5, Awl increased for SiC9 and SiC10 with changing

precursor ratio. Cwl was almost contant within 5 mg/min for different crucible shapes

while the large amount of weight gain was observed for SiC10.

Figure 5.1(b) shows that for the arc current of 100 A, the Awl was highest for SiC2

(crucible CR2, precursor Si), further decreased for SiC4 (crucible CR3, precursor Si:C=1:1),

slightly increased for SiC6 (crucible CR4, precursor Si:C = 1:1) and again decreased for

SiC8 (crucible CR5, precursor Si:C = 1:1). Cwl lied within 5 2 mg/min for different

crucible shapes.

For the arc current of 80 A as well as 100 A, it could be observed that Y showed

similar trend as that of Awl (Figure 5.1 (a)-(d)).


136

For crucible CR2 and precursor Si, Y and Awl increased almost three times while Cwl

reduced very slightly by increasing current from 80 A (Figure 5.1 (a) & (c)) to 100 A

(Figure 5.1 (b) & (d)). Y and Awl did not show much change for crucibles CR3, CR4 and CR5

whereas Cwl reduced very slightly by increasing current from 80 A (Figure 5.1 (a) & (c)) to

100 A (Figure 5.1 (b) & (d)). Exception to this was crucible CR5 for which, Cwl showed

negative value for 80 A (Figure 5.1 (a)), which meant that its weight increased. For crucible

CR5, some experiments were carried out by changing precursor atomic ratio at 80 A arc

current (Samples SiC9 and SiC 10). As the atomic ratio of Si:C was varied as 5:5, 6:4 and

7:3, Y and Awl both increased whereas Cwl became more negative.

Y, Awl and Cwl depend on how the energy supplied by the plasma is utilized. If the

energy dissipation at the surface of anode is favored by the geometry of the electrodes then

anode evaporation is more. Here, SiC is a two element system, i.e. Si and C. Energy of

plasma is utilized for evaporation of Si as well as C and then their reaction to form SiC

(Formation of SiC is an endothermic reaction as is reported by ref [26]). So if Si and C

evaporate without reacting then energy is utilized for evaporation only leading to increased

evaporation rate. When the content of SiC in final product is greater, the yield (Y) is

expected to decrease because enthalpy is being taken up for bond formation between Si and

C, eliminating the possibility of formation of un-reacted phases. This is expected to

eliminate un-reacted Si totally for optimized condition which is explained in the next sub-

sections in detail. However, this data can be used to judge the formation of SiC as a

preliminary tool.

5.2.2.2 XRD Analysis

Figure 5.2 shows the XRD patterns of the as synthesized samples. The standard line

patterns for the different crystalline phases are plotted at the bottom section of the figure.

The highest intensity peak for Si occurs at 28.508° for (111) plane (JCPDS Card No.

772110) whereas for C it occurs at 26.228 for (002) plane (JCPDS Card No 751621). For

SiC, JCPDS Card nos. 742307, 291130, 731749 and 291131 were referred for 3C, 2H, 4H

and 6H polytypes of SiC respectively. As can be observed from the standard line patterns for

different polytypes of SiC in figure 5.2, peak at 35.6° is the highest intensity peak for almost

all the polytypes. Thus, intensity corresponding to this peak was considered for analysis.


137

From figure 5.2, it can be observed that each of the samples consist of SiC along with

impurities either of Si and C or both.

Figure 5.2 X-Ray diffraction patterns of SiC- nanoparticle samples synthesized by thermal plasma.

I. Impurity analysis

Sample SiC1 and SiC2 showed the presence of peak at 2θ = 28.5° (Figure 5.2) which

belong to Si (111) plane and the peak around 35.6° which belong to SiC. The crucible used

for the synthesis of these two samples was the conical cavity crucible CR2 with silicon


138

precursor in it. Although the precursor consisted of silicon alone, still SiC was formed. This

happened due to the erosion of graphite crucible by hydrogen present in argon that increases

the enthalpy [27] of plasma and led to the formation of SiC. The weight percent of SiC in

the synthesized samples increased with the increase in arc current. This happened, because

as the current increases the enthalpy of plasma increases and carbon evaporated from the

crucible increases, that reacts with Si to form SiC. Although the percentage of SiC improved

at arc current of 100 A, the aim was to fully remove the Si impurities and avoid the use of

H2 during synthesis. Thus, further experiments were carried out using Si:C in 1:1 molar ratio

as precursor and two stage cylindrical cavity crucibles were used to reduce the cooling of the

anode.

Sample SiC3 and SiC4, which were synthesized using 3 cm diameter crucible CR3

in presence of Ar, consisted of Si as well as C impurities as can be observed from the peak

for graphitic structure around 2θ = 26.2° and silicon around 28.5° along with the peak for

SiC around 35.6°. Also, it can be observed that the silicon impurities increased at arc current

of 100 A (SiC4) than at 80 A. Thus, it was observed that the heat flux of plasma was not

sufficient for equal evaporation and subsequent reaction of silicon and carbon together. The

melting and boiling point of Si are 1404°C and 3227°C respectively at atmospheric pressure,

while graphite directly sublimate at 3798°C. Also, the thermal conductivity of graphite is

greater than that of silicon. Hence as soon as graphite receives heat, it transfers to silicon;

resulting in to the evaporation of Si.

The diameter of crucible was reduced to 1.7 cm (CR3) to increase the heat flux, thus

the bombardment of the electrons on the surface of the precursor which could sublimate

carbon and allow the reaction of silicon and carbon. As a result, the content of silicon

impurity was seen to be decreased and graphite impurities to be increased in SiC (Sample

SiC5 and SiC6 in figure 5.2) in comparison to Samples SiC3 and SiC4 respectively. This

might be ascribed to the reason that the carbon atoms evaporating from the walls of crucible

condense outside the plasma plume. Here, silicon atoms are not present, so carbon atoms

don’t interact to form SiC. Unlike samples SiC3 and SiC4, the intensity of XRD peaks of

silicon decreased in SiC6 (100 A) than SiC5 (80A) on increasing current. Still the peak for

silicon could be observed for this crucible CR3.


139

So the crucible diameter was further reduced to 1 cm (CR5). The samples SiC7 and

SiC8 were synthesized using CR5 at arc current of 80 A and 100 A respectively. From the

X-Ray diffraction peaks for C, Si and SiC (Figure 5.2) it can be observed that at 80 A

(sample SiC7) the highest intensity peak for Si at 2θ = 28.5° along with all other peaks for Si

were absent. So, the motive was achieved in sample SiC7. Still the synthesis was carried out

at 100 A (SiC8) with a view to increase yield but again the Si-impurities were observed in

XRD.

Thus, for two stage cylindrical cavity crucibles it can be observed that as the

diameter of the anode crucible is reduced the percentage of Si - impurity decreases, while

the percentage of C - impurity increases. In case of 3 cm diameter (CR3), due to the

fluctuations in plasma, silicon and carbon condense without reacting with each other.

Nevertheless, when the diameter is reduced, the fluctuations in plasma are in a smaller area.

So, the heat required for reaction is available that results in the increase of silicon carbide

percentage.

In sample SiC7 (CR5, 80 A and Si:C precursor ratio of 1:1) as discussed above Si

free product was formed but consisted of C- impurities. Hence, the Si:C ratio was increased

to 4:3 (SiC9) and 7:3 (SiC10) to reduce carbon impurities. From the XRD peaks (Figure 5.2)

it is seen that Si impurity was absent but C impurities did not show much difference for

SiC9 when compared to SiC7. For SiC10, Si impurity reappeared.

The relative weight percentage of impurities, with respect to SiC, was estimated

from the XRD pattern using the formula [28],

, (5.5)

where, is the weight percentage of single phase material B, and the most intense

peaks for single phase materials A and B respectively and K is constant. The weight percent

obtained by this method consist of an error of ±5% [28].

To determine the value of K, for Si and SiC as well as C and SiC, the commercial

powders of these materials with particle size in the range of 1 to 100 μm were mixed in

known ratios. From the intensity ratios of XRD patterns of these compositions K for these

pairs was calculated using formula given by equation 5.5. The percentage of Si, C and SiC

calculated by this method for all the as-synthesized samples are mentioned in Table 5.2.


140

Table 5.2 wt% of impurities in SiC samples calculated from XRD pattern and TGA

Sample Different components

present (wt %, from XRD)

Ratio of intensity of XRD

peaks

Wt% of

carbon

by TGA SiC Si C I(38.2)/I(35.6) I(33.6)/I(41.4)

SiC1 48 52 - 0.042 2.36 -

SiC2 71 29 - 0.014 0.4 -

SiC3 78 21 0.96 0.06 2.1 -

SiC4 71.5 26.2 2.3 0.09 1.67 5%

SiC5 90.2 8.2 1.6 0.144 2.96 25%

SiC6 93 5.4 1.6 0.0713 1.86 22%

SiC7 98.2 0 1.8 0.133 2.06 10.1%

SiC8 93.5 5 1.5 0.07 1.92 30%

SiC9 98.0 0 2 0.133 3.24 18%

SiC10 98 1.4 0.6 0.06 1.20 10.2%

II. Polytypes of SiC

The above discussion was limited to the presence of Si and C – impurities in the SiC

samples. Now, we discuss about the polytype formation of silicon carbide. Different

polytypes and their crystal structures have already been discussed in chapter 1, section

1.5.2.1. As mentioned by Seo et al. [29], it is not very simple to determine the amount of

different polytypes present in the samples because of the ambiguity always caused by the

differences, for example, the degree of crystallinity, particle size, stacking fault density

among the polytypes involved.

Polytypes of SiC arise from different periodic stacking sequences of bilayers of Si

and C. The stacking sequence does not significantly alter bond-lengths or affect bulk

density. With Si-C bond-length of 1.89 Å, bilayers are spaced 2.52 Å apart. If the number of

bilayers in the unit cell is even the symmetry is hexagonal otherwise cubic or rombohedral.

Thus, XRD peak corresponding to bilayer-spacing of 2.52 Å is present in almost all the

polytypes. Along with this peak several other peaks are also common in the polytypes as can

be observed from figure 5.2, where the XRD pattern for different polytypes of SiC (JCPDS

Card No. 74-2307 for 3C-SiC, JCPDS Card no. 29-1131 for 6H-SiC, JCPDS Card no. 73-

1749 for 4H-SiC and JCPDS Card no. 29-1130 for 2H-SiC) are plotted. Many researchers


141

have reported that the diffraction peak appearing at 33.6° is due to stacking faults found in

β-SiC and the ratio of the intensity of peaks at 33.6° and 41.4° quantify the stacking faults

[30]. However, this peak corresponds to the stacking faults alone only when peak at 38.2°,

which belong to α-SiC explicitly, is absent [29]. Here, in the present case both the peaks at

33.6° and 38.2° are present. Thus, exact calculation of the phase ratio could not be done, but

it could be observed that both α and β phases of SiC are present in the samples. However, a

comparative intensity ratio of peaks at 38.2° and 35.6° can be used to observe the

comparative presence of phases in the samples.

The ratio of intensity of diffraction peaks at 38.2° and 35.6° as well as 33.6° and

41.4° are listed in table 5.2. From these values, it is observed that for each sample, the

stacking fault decreases at an arc current of 100 A as compared to that of 80 A. From the

values listed in table 5.2, it can be observed that the phase ratio does not follow specific

trend. The exact details about the ratios of different α-polytypes were also difficult to

calculate.

5.2.2.3 Thermogravimetry analysis (TGA)

To determine the content of C and Si impurities in SiC, TGA was carried out for all

samples. Figure 5.3 shows the TG plots for as synthesized SiC samples.

Figure 5.3 Thermogravimetric graphs of all as synthesized SiC samples.

The weight loss from room temperature to about 200°C is due to desorption of

adsorbed gases and moisture. Oxidation of amorphous carbon nanostructures start from


142

500°C [31] while the oxidation of SiC nanoparticles, start from about 750°C [32]. Weight

gain from room temperature is due to presence of silicon nanoparticles that oxidize right

from room temperature. However, after a passivation layer is formed on SiNPs further

oxidation is stopped. On further increase in temperature, when the oxide layer breaks,

oxidation once again begins, like has been observed in the samples SiC1, SiC2 and SiC3,

which contain higher quantity of silicon.

The sample SiC3 shows the total rise phenomenon with some plateau regions in the

range 56°C to 95°C, 230°C to 340°C and 560°C to 600°C. The sample consists of silicon

(observed from XRD), so it shows weight gain. However, these plateau regions might be

due to compensation of weight by removal of adsorbed gases and the amorphous carbon

content present in the sample or due to oxide passivation of SiNPs. The sample SiC4 shows

weight loss right from the beginning.

The increased carbon content can be clearly observed as a steep weight loss of 5 %

starting after 520°C till 650°C indicating the content of the carbon present in the sample.

Sample SiC5, SiC6, SiC7, SiC8, SiC9 and SiC10 show weight loss of 25%, 22%, 10.1%,

30%, 18% and 10.2% respectively staring after 520°C and ending at different temperatures

between 660°C to 725°C due to loss of carbon as CO2. It is noticeable that the sample that

shows a higher percentage of carbon content shows a higher weight loss in the region below

500°C. This shows that amorphous and crystalline carbon nanostructures adsorb moisture

and other gases that get released due to heat resulting in to the weight loss. It was also

observed that the samples SiC7 and SiC9 do not show weight gain except above 700°C

which is due to oxidation of SiC indicating absence of any trace of Si in the samples.

Content of Si could not be calculated exactly but content of carbon could be

calculated considering an error of 2%. The values of carbon content calculated from TGA

for all samples are also mentioned in table 5.2.

5.2.2.4 Microstructure analysis using transmission electron microscope

TEM and HRTEM of all the samples were carried out in order to observe the

microstructure of SiC samples. All the samples consisted of hexagonal, triangular and

truncated-triangular two-dimensional sheet-like as well as pyramidal geometry that belonged


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to SiC (Figure 5.4, 5.5, 5.6, 5.7 and 5.8). SiNSs were present in the form of spherical

particles and nanowires (observed very rarely) in the samples (Figure 5.7).

Figure 5.4 (a) and (b) show the TEM images for the samples SiC1 and SiC2

respectively. The spherical (belonging to Si), triangular (belonging to SiC) and sheet like

(belonging to SiC) particles could be observed in both the samples.

Figure 5.4 TEM micrographs of samples (a) SiC1, and (b) SiC2.

Figure 5.5 TEM micrographs of samples (a) SiC3, and (b) SiC4.

Figure 5.5 (a) and (b) shows the TEM micrographs for the samples SiC3 and SiC4

respectively. As was observed in XRD, spherical nanoparticles (Si nanostructures) are

observed to increase and the triangular particles are observed to decrease in sample SiC4 as


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compared to SiC3. The samples discussed so far up to SiC4 had Si-impurities as a major

content. The samples SiC5 and SiC6 consisted of similar features as that of SiC3 and SiC4,

however, there was presence of C-impurities also alongwith that of silicon.

Figure 5.6 TEM micrographs of (a) carbon hollow and graphene like structures, and (b) graphitic

nanostructures.

As was observed in XRD, the sample SiC7 showed formation of Si-free SiC.

However, the samples synthesized further i.e. SiC8-SiC10 showed the presence of Si-

impurities except SiC9. The carbon impurities consisted of carbon nanostructures in the

form of fullerene like hollow structures (Figure 5.6 (a)), few layered graphene like sheets

and graphitic structures (Figure 5.6 (b)). The HRTEM image of graphitic structure is shown

in the inset of Figure 5.6 (b). The lattice spacing of 3.4 Å belongs to (002) peak of

hexagonal graphite.

Figure 5.7 (a) and (b) shows the TEM micrographs of as synthesized samples SiC7

and SiC9 respectively. The insets show the selective area electron diffraction (SAED)

patterns which shows the rings corresponding to (111), (200), (220) and (311) planes of β-

SiC and (002) plane of graphite. Thus, all the particles belonged either to SiC or C. These

samples were calcined at a temperature of 750°C for 20 min in air to remove the graphitic

particles. This temperature was achieved starting from room temperature in 120 min. After

20 min of calcination at 750°C it was allowed to cool by itself. Figure 5.7 (c) and (d) shows

the TEM image of the sintered samples SiC7 and SiC9 respectively along with the SAED


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patterns (inset) that show diffraction rings corresponding to β-SiC only. This indicates that

the carbon impurities could be removed completely after heat treatment.

Figure 5.7 TEM micrographs of (a) as synthesized sample SiC7, (b) as synthesized sample SiC9, (c)

heat treated sample SiC7, and (d) heat treated sample SiC9 (Insets show the SAED patters of the

corresponding samples).

Overall, based on the analysis of TEM images, following observations could be summarized

as,

a. The average size of the particles reduced from 60 nm to 20 nm.

b. SiC-sheet like structures were more common in anode crucible with 3 cm diameter

(Figure 5.4 (a) and (b)).

c. With the reducing size of anode crucible, the pyramidal structures and three

dimensional structures became more prominent.

Apart from these, the study of growth directions preferred by SiC for the formation

of different microstructures was also interesting. These were studied by HRTEM. The Si-C


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bilayer spacing in SiC is 2.52 Å, which is common in all polytypes of SiC irrespective of the

stacking sequence [33,34]. In this scenario, very well defined HRTEM images or SAED

pattern with particular zone axis observed on a single particle can only give exact

information about the polytype. Hexagonal and cubic phases can be differentiated by

observing the lattice spacing of the particles. If the lattice spacing of 2.52 Å is solely

observed without any other lattice spacing there are fair chances that the particle belongs to

cubic phase. But the presence of any lattice spacing of about 2.6 Å and 2.4 Å indicates the

stacking sequence other than ABC i.e. cubic phase (β-SiC) or presence of 2.6 Å lattice

spacing along-with 2.5 Å lattice spacing alone can be regarded to stacking fault in β-SiC.

Although XRD data and SAED patterns show the presence of β-SiC as a major part

of the sample, HRTEM analysis of some typical faceted structures show the signature of α-

SiC only, discussion about this will follow.

Figure 5.8 (a) TEM micrograph of SiC sample showing typical faceted structures (b) HRTEM image

of the hexagonal 2D structure which is further magnified in (c) with its FFT image in (d), (e) TEM

micrograph of single hexagonal structure with corresponding SAED pattern in inset, (f) schematic

showing possible growth direction resulting in the formation of hexagonal 2D structure, (g)

schematic showing possible growth direction resulting in the formation of triangular 2D structure,

(h) schematic showing possible growth direction resulting in the formation of triangular pyramidal

structure.


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The truncated hexagonal sheets belonged to α-SiC. Figure 5.8 (b) shows the HRTEM

image of one such sheet from sample SiC1. The lattice spacing is found to be 2.66 Å

corresponding to (100) plane in three directions. The angle between two lattice spacings was

found to be 60°. The Fast Fourrier Transform (FFT) of the image (Figure 5.8 (d)) shows the

hexagonal symmetry and the spots correspond to α-SiC. Figure 5.8 (e) shows the hexagonal

sheet from sample SiC7 and the corresponding SAED pattern in the inset. The SAED

patterns show hexagonal symmetry and spots have spacing of 2.66 Å for inner hexagonal

spots marked by first ring from the centre, 1.55 Å (marked by second ring from the centre)

for second hexagonal spots and 1.34 Å (marked by third ring from the centre ring 3) for

outermost hexagon. From the SAED patterns it is clear that the sheet belongs to α-SiC, still

the hexagonal polytype could not be confirmed. But it could be concluded from HRTEM

images that the hexagonal sheet grows perpendicular to c-axis, in the directions < ,

< >, < >, < >, < > and < > as shown in figure 5.8 (f). Some

triangular sheets were also found which show similar lattice spacing and SAED pattern as

that of hexagonal sheets. The growth directions are shown in figure 5.8 (g) which is again

perpendicular to c-axis along < >, < and . Some triangular pyramidal

structures were also observed (Figures 5.8(a)) that again also belonged to α-SiC. The growth

directions of such structure, is proposed in figure 5.8 (h). In this case the growth might be

taking place along four directions which consists of planes , , < >

and < > where k can be any integer constant.


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Figure 5.9 TEM micrograph showing different structures of SiC nanoparticles. Insets 1, 2, 3 and 4

show FFT from the region marked by square 1, 2, 3 and 4.

Figure 5.9 shows some triangular and other structures along with the images of Fast

Fourier transform (FFT) taken on the squares marked on the images. Square 1, 2 and 4 show

the presence of lattice spacings of 2.6 Å, 2.4 Å as well as 2.5 Å showing the clear presence

of α-SiC. But, square 3 shows the presence of lattice spacing of 2.5 Å only, thus it belongs

to β-SiC. Such observations were found on some other particles also.

Figure 5.10 (a) TEM micrograph of triangular shaped SiC nanoparticles, (b) TEM micrograph of

same triangular shaped SiC nanoparticles from different view, (c) TEM micrograph of a structure

observed in SiC samples, (Upper insets show HRTEM images of red squares and lower insets show

the corresponding FFT image).

Figure 5.10 (a) shows the image of triangular particle of SiC which shows a lattice

spacing of 2.5 Å throughout the particle (Figure 5.10(a); upper corner inset shows processed


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HRTEM image of the area marked by a red square and lower corner shows its FFT image).

Same particles, when observed with different angle show lattice spacing of 2.5 Å as well as

2.6 Å (Figure 5.10 (b)). The lattice spacing of 2.6 Å here can be due to stacking fault in β-

SiC as the angle between these two spacing is very small. Figure 5.10 (c) also shows the

rectangular region that consists of lattice spacing of 2.5 Å only. Thus, these structures might

belong to β-SiC, still exact confirmation could not be made. Samples consisting of silicon

impurities consisted of the few instances of Si-SiC nanojunction. The spherical particle

attached to the rectangular part (Figure 5.10 (c)) belongs to cubic silicon.

Figure 5.11 (a) TEM micrograph of SiC-Si nanojunction formation (lower hexagonal sheet belongs

to SiC while the hemispherical structure belongs to Si), (b1) and (b3) show the magnified images of

square 1 and 2 in (b) and (b2) and (b4) show the corresponding FFT images. (c) TEM micrograph

consisting of Si and SiC junction, (c1) FFT of square 1in (c) showing presence hexagonal Si, and

(c2) FFT of square 2 in (c) showing presence of hexagonal silicon carbide.

Many such nanojunctions were observed on hexagonal and truncated-triangular SiC

sheet structure (shown in figure 5.11(a), (b) and (c)). Figure 5.11 (b1) and (b3) show

HRTEM images of the areas marked by square 1 and 2 respectively while figure 5.11 (b2)

and (b4) show their FFT images. The lower hexagonal structure shows lattice spacing of 2.5


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Å and 2.6 Å with an angle of 60°. This angle of 60° is reported for 6H-SiC [35]. The

hemispherical structure shows the lattice spacing of 3.1 Å corresponding to (111) plane of

cubic Si. The lattice spacing of 5.3 Å can be ascribed to Moiree fringes. At some places

where Si has grown on hexagonal SiC, Si exhibits hexagonal crystal structure. One such

image is shown in figure 5.11 (c). The FFT of square 1 and 2 in figure 5.11 (c) are shown in

figure 5.11 (c1) and (c2) respectively. The FFT shows hexagonal symmetry with lattice

spacing of 3.3 Å with an angle of 60° between the two planes [36], [37]. Such a feature

cannot be observed in cubic silicon on observation through any zone axis. Thus, it confirms

the growth of hexagonal of Si on hexagonal SiC.

Thus, from the overall TEM analysis it can be confirmed that β-SiC is not present in

the form of sharp edged two dimensional structures. It is present in the form of smaller and

triangular particles as can be observed in figure 5.9 and 5.10.

5.2.2.5 UV-Visible spectroscopy analysis

Figure 5.12 (a) shows the UV-Visible spectra of samples SiC1 and SiC2 which

consist of Si and SiC nanoparticles.

Figure 5.12 UV-Visible absorption spectra of samples SiC1 and SiC2.

The spectra show the absorption edge at 540 nm and 533 nm for samples SiC1 and

SiC2 respectively. These wavelengths correspond to energy of 2.29 eV and 2.32 eV

respectively. These values are less than the band gap for bulk β-SiC which is 2.39 eV. This

may be due to overlap of absorption from Si which is also observed throughout the

remaining spectra.


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The samples consisting of Si as well as C impurities showed nearly simillar spectra

of absorption in the UV-Visible region. Figure 5.13 (a) shows the UV-Vis spectra of sample

SiC3 which shows the increases in absorption up to 400 nm, then shows a bend (as was

observed in samples SIC1 and SiC2) and then remains almost constant throughout the

visible a region. This continuous absorption is because of the presence of Si and C

impurities which absorb the entire photons in the visible region.

Figure 5.13 (a) UV-Visible absorption spectra of sample SiC3, and (b) samples SiC7 and SiC9.

Figure 5.13 (b) shows the UV-Vis spectra of samples SiC7 and SiC9 which consisted

of C impurity only. The spectra shows continuously increasing absorption in the visible

region and it is very difficult to conclude anything about the bandgap of SiC. Thus, the

spectra were recorded after calcinations of these samples.

Figure 5.14 shows the spectra of the samples SiC7 and SiC9 after calcination. The

spectra shows the presence of two band edges of which one is present at 326 nm

corresponding to 3.8 eV (about 3.4eV for bulk α-SiC) for both the samples. The other band

edge position is different for the two samples; for SiC7 it is at 432 nm corresponding to 2.8

eV and for SiC9 it is at 456 nm equivalent to 2.7 eV. The band edge is not very sharp owing

to particle size distribution. Both the band gap values are greater than the values for the bulk

α-SiC and β-SiC samples. This might be the result of quantum confinement effect arising

from small size of SiCNPs. Thus, it is clear that the samples consist of mixed polytypes of

SiC, but the exact ratio could not be quantified on the basis of available evidences.


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Figure 5.14 UV-Visible absorption spectra of samples SiC7 and SiC9 after calcination.

5.2.3 Conclusions

The DC direct arc thermal plasma is an effective tool for the synthesis of SiCNPs.

The heat dynamics of arc plasma strongly depends on the geometry of the electrodes. Also,

the evaporation of the anode and the reaction occurring near the anode are dependent on

crucible geometry. When the diameter of anode was comparable with that of cathode, the

stability of arc was found maximum. The parameters of synthesis were successfully

optimized for the synthesis of silicon free SiC nanoparticles. This was observed using XRD,

TGA, TEM and UV-Vis Spectroscopy. The results showed that the sample synthesized

using 1cm (diameter) double stage cylindrical crucible and 9cm (diameter) graphite cathode

at an arc current of 80 A and arc voltage of 12-14 V in presence of 500 torr Ar yielded the

SiC sample free from Si-impurity. Although Si impurities were successfully avoided the

final product of SiC consisted of mixed polytype system. This was observed from the XRD,

TEM and UV-Vis Spectra. While observing morphology in TEM different microstructures

were seen. Detail HRTEM and SAED analysis showed that the structures with sharp edges

belonged to α-SiC and their possible growth directions are proposed based on these studies.

The three dimensional triangular structures and few of the small spherical structures

belonged to β-SiC. UV-Vis spectra also show the band edge for both α-SiC and β-SiC

nanostructures. However, the exact polytype ratio could not be obtained.


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5.3 SiCNPs - diglycidyl ether bisphenol A (DGEBA) epoxy polymer

composites

As discussed earlier DGEBA possess many favourable properties which make it

useable for numerous applications. But, it suffers from the major drawback in terms of poor

resistance to crack initiation. SiC is a hard material, so making composite of such two

materials would be helpful for fulfilling the requirement. In this part, a brief background

about epoxy especially DGEBA and processes associated with it are discussed. Further,

experimental procedure adopted for the preparation of composites and discussion about the

results has been included.

5.3.1 Epoxy polymers

Epoxy resins represent an important class of polymers primarily due to their

versatility. High degree of crosslinking and the nature of the interchain bonds give cured

epoxies many desirable characteristics. These characteristics include excellent adhesion to

many substrates, high strength (tensile, compressive and flexural), chemical resistance,

fatigue resistance, corrosion resistance and electrical resistance. In addition, processing is

simplified by low shrinkage and lack of volatile by-products. Properties of the uncured

epoxy resins such as viscosity, which are important in processing as well as final properties

of cured epoxies such as strength or electrical resistance, can be optimized by appropriate

selection of the epoxy monomer and the curing agent or catalyst. Because of the ease of

application and desirable properties, epoxies are widely used for coatings, corrosion

protectants, electric encapsulants, fiber optic sheathing, flooring and adhesives.

5.3.2 Diglycidyl ether bisphenol A (DGEBA)

Diglycidyl Ether Bisphenol A (DGEBA) type epoxy resin is the most widely used

matrix for innumerable applications, owing to its well balanced chemical, adhesive, thermal

and processing characteristics.

Epoxies are characterized by the presence of one or more epoxide functional groups

on or in the polymer chain. The epoxide group is planar, with a three-membered ring

composed of one oxygen and two carbon atoms as shown in figure 5.15. DGEBA is a


154

derivative of bisphenol A and glycidol marked in figure 5.16. It consists of two epoxide

groups at the ends.

Figure 5.15 Chemical formula of epoxide group.

Figure 5.16 Chemical formula of DGEBA.

5.3.3 Curing of epoxy

The curing reaction of epoxide is the process by which one or more kinds of

reactants, i.e., an epoxide and one or more curing agents with or without the catalysts are

transformed from low molecular weight to a highly crosslinked structure. As mentioned

earlier, the epoxy resin contains one or more 1,2-epoxide groups. Because an oxygen atom

has a high electronegativity, the chemical bonds between oxygen and carbon atoms in the 1,

2-epoxide groups are the polar bonds, in which the oxygen atom becomes partially negative,

whereas the carbon atoms become partially positive. Because the epoxide ring is strained

(unstable), and polar groups (nucleophiles) can attack it, the epoxy group is easily broken. It

can react with both nucleophilic curing reagents and electrophilic curing agents. The curing

reaction is the repeated process of the ring opening reaction of epoxides, adding molecules

and producing a higher molecular weight and finally resulting in a three - dimensional

structure. Figure 5.17 shows how epoxy groups react with amine groups to form crosslinks.

http://en.wikipedia.org/wiki/Bisphenol_A

http://en.wikipedia.org/wiki/Glycidol


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Figure 5.17 Reaction of epoxy group with amine group.

5.3.4 Procedure of preparation of SiC nanoparticles - DGEBA composites

As synthesized SiCNPs were not dispersible directly into DGEBA. Hence, they had

to be dispersed in some solvent prior to dispersion in DGEBA such that DGEBA is also

soluble in that solvent. Hence, initially dispersion studies were carried out in following

solvents; 1. Benzyl alcohol, 2.Benzene, 3.Isopropyl alcohol, 4. Ethanol amine, 5. Toloune, 6.

Chloroform and 7.Ethanol.

20 mg SiC powder was added in 1ml of each of the solvents followed by

ultrasonication for 30 min. The dispersions were then kept still for 24 hours. It was observed

that the particles settled at the bottom of the testtubes. So, now the dispersions were

ultrasonicated in the ultrasonic bath at 65°C for 30 min and again the stability was observed

after 24 hours. Figure 5.18 (a) and (b) show the photographs of different dispersions just

after ultrasonication and 24 hours of ultrasonication. It was observed that the particles could

be uniformly dispersed in benzyl alcohol, isopropyl alcohol, ethanol amine and ethanol. But,

the dispersions were found to be stable in benzyl alcohol and isopropyl alcohol for a longer

period of time. Solubility of epoxy is better in benzyl alcohol, thus it was used for the

preparation of composites.


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Figure 5.18 The photograph of different dispersions just after ultrasonication at 65°C for 30 min and

after 24 hours of ultrasonication. (1.Benzyl alcohol, 2.Benzene, 3.Isopropyl alcohol, 4.Ethanol

amine, 5. Toloune, 6. Chloroform, 7. Ethanol).

For the preparation of composites, 4 gram of DGEBA was used. 6 different

composites with different filler concentrations of 0%, 0.25%, 0.50%, 1%, 1.5% and 2%

were prepared. The amount of SiC powder for filler percentage of 0.25%, 50%, 1%, 1.5%

and 2% of 4 gram of DGEBA came out to be 0 gm, 0.01gm, 0.02gm, 0.04gm, 0.06gm and

0.08gm respectively. This quantity of SiC powder was taken in test tube each containing 1

ml benzyl alcohol in order to disperse powder uniformly. This mixture was sonicated in

ultrasonic bath for 30 min at 65ᴼC. 4 gm of DGEBA was taken in 6 different beakers. The

above dispersions of SiCNPs in benzyl alcohol was added in each beaker with continuos

stirring with the help of a spatula. These mixtures were ultrasonicated at 65°C for 30 min.

Then triethyl tetra amine (TETA) was added into the mixture, which acted as a hardner. The

amount required was 0.5106 gram which was derived from following formula (derived from

the equivalent wt.),

Amount of TETA required= (Weight of DGEBA*24)/(188) (5.6)

This mixture was mixed uniformly with continuos stirring for 20 min. Equal amounts

of mixtures were then poured in 6 different equal sized petri dishes. The mixtures were kept


157

still in levelled dry place and allowed to react for 24 hours. After 24 hours, the samples were

heated at 60°C in oven for complete curing for 30 min and then heated at 180°C to remove

the remnant benzyl alchohol.

5.3.5 Study of properties of SiCNP – DGEBA composites

Figure 5.19 shows the photograph of pure DGEBA (0% filler) and figure 5.20 shows

the photograph of SiCNPs – DGEBA composites with increasing concentration of filler

from left to right 0.25%, 0.50%, 1%, 1.5%, 2% respectively.

Figure 5.19 The photograph of pure DGEBA.

Figure 5.20 The photograph of nano SiC – DGEBA composites with increasing concentration of

filler from left to right (0.25%, 0.50%, 1%, 1.5%, 2% respectively).

5.3.5.1 Study of SiCNP-DGEBA composites by scanning electron microscope (SEM)

The microstructures of composites were observed by SEM. Figure 5.21 shows the

SEM images taken at the cross section of different composites.


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Figure 5.21 SEM images of different composites, (a) 0% filler, (b) 0.25% filler, (c) 0.50% filler, (d)

1% filler, (e) 1.5% filler, (f) 2% filler.

It was observed that the 0% as well as 0.25% filler DGEBA polymer showed clean

surface. The images of other composites show the presence of some particles. They are

observable in SEM indicates that the SiCNPs have agglomerated in the composites. The

agglomerated particle size as observed in SEM varies between 50 nm to 300 nm. The

density of particles is seen to increase with increasing filler concentration.

5.3.5.2 Hardness measurements

The hardness is defined as resistance of a metal to plastic deformation. However, the

term may also refer to stiffness or temper or resistance to scratching, abrasion, or cutting.

The hardness testing of polymers is most commonly measured by the Shore hardness test or

Rockwell hardness test. Both scales provide an empirical hardness value that doesn't

correlate to other properties or fundamental characteristics. Shore hardness, using either the


159

Shore A or Shore D scale, is the preferred method for polymers. The Shore A scale is used

for 'softer' rubbers while the Shore D scale is used for 'harder' ones.

Shore hardness is measured using a durometer that consists of a spring-loaded

indenter mounted with diamond tipped hammer in a graduated glass tube which is allowed

to fall from a known height on the specimen to be tested. The hardness number depends on

the height to which the hammer rebounds; the harder the material, the higher the rebound. If

the indenter completely penetrates the sample, a reading of 0 is obtained, and if no

penetration occurs, a reading of 100 results. The reading is dimensionless. The greater the

number, the greater is the resistance. The results obtained from this test are a useful measure

of relative resistance to indentation of various grades of polymers. The durometer from

Hiroshima (model no. RR-12) was used for the measurements in this work.

Table 5.3 shows the Shore D hardness values of SiCNP-DGEBA composites. It is

observed that with increasing filler concentration the Shore D hardness of the composites

increases. The Shore D hardness for pure DGEBA and composite with 0.25% filler is found

to be 70. It increases gradually with increase in the filler concentration and is found to be

highest for 2% filler concentration i.e. 85. Thus 2% filler concentration changes the property

of DGEBA effectively.

Table 5.3 Hardness values of SiC-Epoxy composites with increasing filler percentage.

Filler percentage (%) Hardness value

0 70

0.25 70

0.50 71

1 72

1.5 75

2 85

5.3.5.3 Chemical stability study

To study the chemical stability of the composites in comparison with pure epoxy,

three different samples with 0%, 1% and 2 % filler concentration were kept in 10 M NaOH,

10 M H2SO4 and NN-dimethylformamide. The weight change in the material was recorded


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after 24 hours of treatment in these solutions and FTIR spectra were recorded to study the

effect of these chemicals on the stability of composites.

Figure 5.22 shows the FTIR spectra of pure DGEBA after treatment with NaOH,

H2SO4 and NN-dimethylformamide. It was observed that NaOH did not have any effect on

pure DGEBA, while due to H2SO4 and NN-dimethylformamide few changes were observed.

By treatment with H2SO4 the peaks at 880 cm-1

, 1105 cm-1

and 1660 cm-1

got modified. The

peak at 880 cm-1

corresponds to stretching mode of C-O-C oxirane group, 1105 cm-1

corresponds to stretching mode of C-O-C of ethers and 1660 cm-1

corresponds to C=C

stretch. Hence, H2SO4 attacks C-O-C bonds creating new C=C bonds. On treatment with

NN-dimethylformamide no change in the spectra was observed, except the appearance of an

extra peak around 1665 cm-1

. This peak can be due to C=C or C=O stretch. Hence, NN-

dimethylformamide is creating these extra bonds on the surface of polymer sample.

Figure 5.22 FTIR Spectra of pure DGEBA after treatment with NaOH, H2SO4 and NN-

dimethylformamide

Figure 5.23 (a) and (b) show the FTIR spectra of SiCNPs-DGEBA composites

treated with H2SO4 and NN-dimethylformamide respectively. No change in the spectra was

observed for different SiCNP concentration, thus due to dispersion of SiCNP no effect in

reaction with H2SO4 and NN-dimethylformamide is observed. This could be owing to the

reason that the particles get submerged within the matrix of DGEBA and hence do not show

difference in surface reactivity of composites.


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Figure 5.23 FTIR Spectra of nano-SiC- DGEBA composites with different filler concentration before

and after treatment with H2SO4

5.3.5.4 Study of thermal properties by using thermo gravimetric analysis (TGA)

In order to study the thermal stability of DGEBA composites TGA was carried out.

Figure 5.24 shows the TGA graphs of SiCNPs - DGEBA composites of different filler

concentration. The weight change was recorded was recorded at the rate of 4°C /min from

room temperature to 500°C in the flow of air at the rate of 20 ml/min. All the curves show

two major weight loses that start at around 150°C and 325°C. These correspond to

evaporation of remaining benzyl alchohol and decomposition of DGEBA polymer. The

percent weight losses of different composites are mentioned in table 5.4. It was observed

that the weight loss for pure epoxy is greater than all the other composites which imply that

the composites have slightly greater thermal stability owing to presence of SiCNPs.

Figure 5.24 TGA graphs of nano-SiC- DGEBA composites of different filler concentration.


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Table 5.4 The percent weight losses of different composites.

Filler Percentage (%) Percent weight loss (%)

0 83.5

0.25 82.2

0.5 82.8

1 80.5

1.5 82.5

2 81.7

5.3.6 Conclusions

In conclusion, the SiCNPs-DGEBA composites are succesfully synthesized. These

NPs could be well dispersed by the use of benzyl alcohol as was easily visible by naked

eyes. However, from SEM micrographs showed the considerable agglomoration in NPs at

higher concentration of SiCNPs. To further improve dispersion, it was required to

functionalize the nanoparticles prior to dispersion. Shore D hardness test has shown that the

hardness increases with increasing filler concentration. The hardness increased from 70 for a

pure epoxy to 85 for 2% filler concentration. These results proved that the addition of

SiCNPs into DGEBA improved the hardness. TGA has also shown slight improvement in

the thermal behaviour of composites with increasing filler concentration.

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Chapter 6

Conclusions and Future Scope

This chapter concludes the findings of the thesis and proposes the work that can be done further in

continuation to the work presented here.

Chapter 6. Conclusions and future scope

166

6.1 Conclusions

The journey of Ph D is a process of learning, grasping and investigating the avenues

necessary for development. One cannot expect a new discovery but yes, completion of one

avenue and setting another on the basis of work carried out! In this process, it would be of

immense pleasure if I was able to unveil a step in a billion.

The thesis consists of a careful experimental piece of work involving vapour phase

synthesis of different forms of nanocrystalline silicon and silicon carbide. Although, all the

results related to the experimental details of the synthesis, characterization and applications

have been discussed in five chapters, it would be noteworthy to highlight the major findings

in the chapter especially devoted for conclusions. In view of complicated procedure during

the vapour phase nucleation and growth, the experimental results seem to be quite important

in the field of semiconductor physics and nanotechnology. Controlling the growth of

nanostructures, is infact, “an art of crystal growth” as has been stated rightly by Gilman.

Such a high temperature synthesis has not been reported very frequently and therefore

demands investigation. It also leads to the possibility of generating, much unrevealed

science behind the mixed nature of sp2 and sp

3 – hybridized structures in Si.

The key finding of thesis is the study of silicon nanostructures, especially nanotubes,

with a point of view of its synthesis, characterization and its possible application. Only

theoretically predicted and experimentally unviable silicon nanotubes could be successfully

synthesized and studied.

The important findings of the thesis are summarized as follows:

1. The parameters to obtain silicon nanotubes were optimized by performing sets of

experiments. The nanotubes could be synthesized in presence of 5 % H2 in Ar (500

torr), at an arc current of 80-90 A and an arc voltage of 12-14 V. The sample of as

synthesized nanotubes consisted of nanotubes and particles in the ratio of ~70:30. The

NPs were spherical in shape with sizes varying between 5-25 nm, while the diameters

of the NTs ranged between 9 nm and 30 nm. A large number of tubes had the diameter

of 14 ± 2 nm with lengths of the order of several hundreds of nm. The nanotubes

consisted of single walled SiNTs, often oxidized, but consisted of some non-oxidized

regions giving the hints about the possibility of presence of sp2 and sp

3 – mixed


167

hybridized silicon. The wall thickness observed in the nanotubes was found to be ~ 0.7

Å. The particles which were present in the sample of SiNTs were also found to be

hollow single walled bucky-ball like spherical structures, similar to those observed in

carbon, in carbon they are formed owing to mixed sp2 and sp

3 – hybridization. This

again hints at the possible presence of mixed sp2 and sp

3 – hybridization obtained in Si.

The experimental observations showed that both, hydrogen and oxygen played

important role in the synthesis of SiNTs.

2. The nanotubes were studied for anti bacterial properties. The anti bacterial activity of

SiNT sample was compared with that of sample consisting of major concentration of

nanoparticles. Four bacteria, two each of Gram positive (Staphylococcus aureus and

Bacillus subtilis) and Gram negative (Escherichia coli and Pseudomonas aeruginosa),

were used for the study. The IC-50 (inhibition of bacteria by 50 %) value for nanotubes

was 200 μg/ml in B. subtilis cultures. 10 μg/ml of nanotubes were proved to be effective

in controlling the S. aureus. With Gram-negative bacteria like E. coli, MIC was found

to be 10 μg/ml for both nanoparticle and nanotube samples. Both the samples were

found to be competent in controlling both Gram-positive and Gram-negative bacterial

strains tested.

3. SiNTs are predicted to be good electron field emitter, so, the field emission properties

of the sample were investigated. A maximum current density of 4.2 mA/cm2

was

attainable at applied electric field of 2.8 V/µm. A low turn on field of merely 1.9 V/µm

was required to draw a current density of 10 µA/cm2. The current stability at 1 µA

preset value is found to be good.

4. The other material, synthesized during this work, was silicon carbide. Synthesis of Si

and C – free SiC nanoparticles was a difficult task and has been worked on since long.

This was achieved by changing the morphology of anode and cathode, and altering the

heat dynamics of the system. The crucible diameter (anode) was changed from 3 cm to

1cm while the geometry was also altered (conical cavity/ cylindrical cavity). When the

diameter of crucible was reduced to 1 cm (cylindrical cavity), the silicon impurity could

be fully avoided during synthesis. Carbon impurities could be further removed by

calcination. Even the particle size distribution could be controlled quite effectively. At

the optimized parameters, the size distribution of the SiC particles was found to be

around 20±15 nm.


168

5. As synthesized SiC nanoparticles consisted of mixed polytype system. TEM

micrographs showed the presence different shapes of SiC nanoparticles. They were

extensively studied by TEM in order to investigate the polytype of nanoparticles with

particular morphology. The growth directions and formation of preferred morphology

for particular polytype has been predicted on the basis of HRTEM and SAED patterns.

6. SiC nanoparticles – DGEBA epoxy composites were fabricated with different filler

concentration. A considerable increase in Shore D hardness is observed for 2 % of

SiCNPs in DGEBA.

6.2 Future Scope

Research never ends, so there is always a scope ahead in any piece of work. Although

the nanotubes were synthesized successfully in this work, there lies a lot of scope in this

field i.e. reducing of oxygen content from nanotubes and opening of these nanotubes to form

single atomic layer sheets i.e. silicene and the study of the band properties of these

structures. These nanotubes can be employed for further applications in devices such as

Lithium ion batteries, in anti-bacterial coatings and composites with polymers, etc.

The impurity free SiC nanoparticles, synthesized in this work, are suitable for

applications like polymer composites, abrasive coatings etc. However, specified applications

like optoelectronics, microwave absorption require phase purity of nanoparticles. Further

modifications are required to control the thermodynamical properties of plasma to achieve

phase purity in SiC. The application of external electric and magnetic field may serve as the

possible solutions in nearby future.

Publications

169

Publications

International Journal publications

1 Arc plasma synthesized Si nanotubes: A promising low turn on field emission

source.

Padmashree D. Joshia)

, Chiti M. Tanka)

, Shalaka A. Kamble, Dilip S. Joag, Sudha V.

Bhoraskar and Vilas L. Mathe. J. Vac. Sci. Technol. B 33(2), Mar/Apr 2015 (accepted)

2 Antimicrobial activity of silica coated silicon nano-tubes (SCSNT) and silica coated

silicon nano-particles (SCSNP) synthesized by gas phase condensation

Chiti Tank, Sujatha Raman, Sujoy Karan, Suresh Gosavi, Niranjan P. Lalla, Vasant

Sathe, Richard Berndt, W. N. Gade, S. V. Bhoraskar, Vikas L. Mathe. J Mater Sci

Mater Med. 24 (2013)1483-90, doi: 10.1007/s10856-013-4896-3.

3 Si nanotubes and nanospheres with two-dimensional polycrystalline walls†

Paola Castrucci‡*a, Marco Diociaiuti‡

b, Chiti Manohar Tank‡

c, Stefano Casciardi,

Francesca Tombolini, Manuela Scarselli, Maurizio De Crescenzi, Vikas Laxman Mathe

and Sudha Vasant Bhoraskar. Nanoscale 4, (2012) 5195 doi: 10.1039/c2nr30910f.

4 Thermal Plasma Assisted Synthesis of Nanocrystalline Silicon—A Review

S. V. Bhoraskar*, C. M. Tank and V. L. Mathe, Nanosci Nanotech Lett. 4, (2012) 1

291-308(18), doi: 10.1166/nnl.2012.1319.

Conference Proceedings

5 Synthesis of Silicon Nanostructures Using DC-Arc Thermal Plasma: Effect of

Ambient Hydrogen on Morphology

Chiti M. Tank, Vijaykumar B. Varma, Sudha V. Bhoraskar, Vikas L. Mathe *

Advanced Materials Research 938, (2014) 76-81

6 Synthesis of silicon nanotubes by DC arc plasma method

C. M. Tank, S. V. Bhoraskar, and V. L. Mathe, AIP Conf. Proc. 1447, (2012) 423

doi:10.1063/1.4710060

Publications

170

Journal publications other than thesis

7 Electric field enhanced photocatalytic properties of TiO2 nanoparticles immobilized

in porous silicon template.

C. M. Tank, Y.S. Sakhare, N.S. Kanhe, A.B. Nawale, A.K. Das, S.V. Bhoraskar, V.L.

Mathe,* Solid State Sciences 13 (2011) 1500 doi:

10.1016/j.solidstatesciences.2011.05.010

8 ECR plasma assisted deposition of nano-TiO2 for repeated applications of

photocatalytic degradation.

Avinash S Bansode, Chiti M Tank, K R Patil, S V Bhoraskar and V L Mathe

Archives of Applied Science Research, 2 (2010) 288

Patents

Plasma-based method for synthesis of nano-sized Silicon carbide

Chiti Tank, S. V. Bhoraskar and V. L. Mathe (Indian Patent under process)

Conferences and Awards

171

Conferences/Symposium/Schools Attended

Poster presentation at Second Conference on Nanotechnology for Biological and

Biomedical Applications (Nano-Bio-Med 2013), held at ICTP, Trieste, Italy during

October 14 – 18, 2013.

Oral Presentation at National Symposium on Emerging Plasma techniques for

Material Processing and Industrial Applications held at Department of Physics,

University of Pune, during Feb, 13-15, 2014.

Oral Presentation at National Conference on Functional Nanomaterials-2013 held

at Department of Physics, University of Pune, Pune during January 31- February 1,

2013.

Poster Presentation in 56th

DAE Solid State Physics Symposium held at SRM

University, Kattankulathur, Tamilnadu during December 19-23, 2011.

Secretary, Raman Memorial Conference-2012 held at Department of Physics,

University of Pune, Pune, during February 22-23, 2012.

SERC School on Nano Optics held at NIT Hamirpur, Himachal Pradesh, India,

during September 13- October 01, 2010.

Awards

Best Oral Presentation award, for paper “Synthesis of SiC nanoparticles by gas

phase condensation using DC-arc thermal plasma.”

National Conference on Advances in Plasma Science and Technology held at Sri

Shakthi Institute Engineering and Technology, Coimbatore during February 19-21,

2015

Best Poster Award, for paper “Synthesis of Silicon Nanostructures using DC- Arc

Thermal Plasma: Effect of Ambient Hydrogen on Morphology” International Conference On Nano Materials: Science, Technology And Applications

(ICNM' 13) held at B S Abdur Rahman University, Vandalur, Chennai -48, Tamil

Nadu, India during December 05 - 07th

2013.

Best Oral Presentation award, for paper “Synthesis of thin walled Silica-coated

Silicon Nanotubes and their Antibacterial Study”

International Conference on Applications of Advanced Materials for Sustainable

Development, held at Nagpur during Jan 17-18, 2014.

Best Oral Presentation award, for paper “Synthesis of SiC Nanosheets and SiC-Si

Nanojunction: Transmission electron Microscopy study”

Raman Memorial Conference-2013 held at Department of Physics, University of

Pune, Pune during February 22-23, 2013.

Documents

List of symbols and abbreviations - Shodhgangashodhganga.inflibnet.ac.in/bitstream/10603/90094/15/15_list_of_publications.pdf2.2.2 Applications of SiC nanostructures 59 Bibliography