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ANDREY OREKHOV

ELECTRON MICROSCOPY STUDY OF STRUCTURALPECULIARITIES OF CARBON MATERIALS


This work presents results of high resolution transmission electron microscopy analysis

of structural peculiarities of some of nanocarbons materials forms. It was

performed the optimization of optical system of aberration corrected transmission electron

microscope to increase of the signal-to-noise ratio. The optimized TEM instruments

were used in this work for structural characterization of the nanodiamonds, two-

dimensional (2D) structures and needle-like diamonds, onion-like nanocarbons and other graphene-based structures, including

composites consisting of linear (1D) CuCl and Hg2Cl2 crystals encapsulated in single-walled carbon nanotubes. Obtained results allowed

appropriate development of production processes for these carbon nanostructures and

understanding of their physical properties.

ANDREY OREKHOV

ANDREY OREKHOV

Electron Microscopy Study

of Structural Peculiarities of

Carbon Materials

Publications of the University of Eastern Finland


Number 258

Academic Dissertation

To be presented by permission of the Faculty of Science and Forestry for public

examination in the Auditorium AU100 in Aurora Building at the University of

Eastern Finland, Joensuu, on December, 22, 2016, at 12 o’clock noon.

Department of Physics and Mathematics

Grano Oy

Joensuu, 2016

Editors: Prof. Pertti Pasanen,

Prof. Pekka Toivanen, Prof. Jukka Tuomela, and Prof. Matti Vornanen

Distribution:

University of Eastern Finland Library / Sales of publications

P.O.Box 107, FI‐80101 Joensuu, Finland

tel. +358‐50‐3058396

www.uef.fi/kirjasto

ISBN: 978‐952‐61‐2381‐3 (Print)

ISBN: 978‐952‐61‐2382‐0 (PDF)

ISSNL: 1798‐5668

ISSN: 1798‐5668

ISSN: 1798‐5676 (PDF)

Author’s address: University of Eastern Finland


P.O.Box 111

80101 JOENSUU

FINLAND

email: [email protected]

Supervisors: Professor Alexander Obraztsov, Ph.D.

University of Eastern Finland


P.O.Box 111

80101 JOENSUU

FINLAND


Professor Yuri Svirko, Ph.D.

University of Eastern Finland


P.O.Box 111

80101 JOENSUU

FINLAND


Reviewers: Professor Angela Vella, Ph.D

Normandie Université, Université‐INSA de Rouen

Groupe de Physique des Matériaux

Avenue de lʹUniversité BP 12

76801 SAINT ETIENNE DU ROUVRAY

FRANCE

email: angela.vella@univ‐rouen.fr

Professor Rupert Schreiner, Ph.D

OTH Regensburg Fakultät AM

Seybothstr. 2

93053 REGENSBURG

GERMANY

email: rupert.schreiner@oth‐regensburg.de

Opponent: Professor Albert Nasibulin, Ph.D.

Skolkovo Institute of Science and Technology

Skolkovo Innovation Center, Building 3

143026 MOSCOW

RUSSIA


ABSTRACT

Carbon materials attract growing interest of researchers due to

their numerous unique properties and potential applications.

These properties are originated of atomic structure of carbon

which exhibits great variety of its nanostructured forms

including graphene, carbon nanotubes, fullerenes,

nanodiamonds and other. Revealing structural peculiarities of

these nanostructured forms of carbon materials is very

important for understanding of origin of their electronic, optical,

thermo‐physical, mechanical and other properties. In this

Thesis, the high resolution transmission electron microscopy is

applied to study structural peculiarities of some of nanocarbon

materials. The Thesis includes brief overview of some forms of

nanostructural carbon materials, description of methodological

problems of the transmission electron microscopy and

possibilities for their resolving. These descriptions are

illustrated by examples of the researches performed by the

author and include his original publications.

Universal Decimal Classification: 538.911, 548.1

CAB Thesaurus: high resolution transmission electron microscopy, analytical

methods, atomic structure; nanostructured materials, nanocarbon materials,

carbon nanotubes, nanodiamonds;

Preface

I would like to express my thanks to all those who have

helped me in this work:

Above all, I am profoundly indebted to my supervisors

Professor Alexander Obraztsov for his guidance, fundamental

insight, his time and great efforts supporting this work. I am

also very grateful to my supervisor Professor Yuri Svirko and to

the Head of the Department of Physics and Mathematics

Professor Timo Jääskeläinen for their support and for providing

me this opportunity to realize my research project in the

University of Eastern Finland.

Author of this Thesis greatly appreciates support from the

collaborators and colleagues from research groups headed by

Professor Alexander Obraztsov, Professor Elena Obraztsova and

Professor Alexander Okotrub which made great and valuable

inputs providing ideas, materials and discussion.

Special thanks are to Professor Andrey Chuvilin for giving

the opportunity to carry out the TEM measurements and

support in image processing and for discussions of results.

I am grateful to the Examiners of this work, Professor Angela

Vella and Professor Rupert Schreiner, and to the Opponent,

Professor Albert Nasibulin, for their review and encouraging

comments.

And finally, I would like express my gratitude to my family,

parents and brothers for their patience, love and comprehensive

support.

This work made with partial support from Russian Science

Foundation (grant #14‐12‐00511).

Joensuu December 22, 2016 Andrey Orekhov

LIST OF ORIGINAL PUBLICATIONS This thesis consists of the review of author’s work in field of

electron microscopy study of structural peculiarities of carbon

materials and the following selection of the author’s

publications, referred to by the Roman numerals I‐VIII:

I D’yakova Yu. A., Suvorova E. I., Orekhov Andrei S.,

Orekhov Anton S., Alekseev A. S., Gainutdinov R. V.,

Klechkovskaya V. V., Tereschenko E. Yu., Tkachenko N. V.,

Lemmetyinen H., Feigin L. A., Kovalchuk M. V. Study of

Structural Order in Porphyrin–Fullerene Dyad ZnDHD6ee

Monolayers by Electron Diffraction and Atomic Force

Microscopy, Crystallography Reports 58(6): 927–933, 2013 DOI:

10.1134/S1063774513060096

II Rybkovskiy D. V., Arutyunyan N. R., Orekhov A. S.,

Gromchenko I. A., Vorobiev I. V., Osadchy A. V., Salaev E.

Yu., Baykara T. K., Allakhverdiev K. R., Obraztsova E. D.

Size‐induced effects in gallium selenide electronic structure:

The influence of interlayer interactions, Physical Review B 84:

085314, 2011 DOI: 10.1103/PhysRevB.84.085314

III Orekhov A.S., Savilov S.V., Zakharov V.N., Yatsenko A.V.,

Aslanov L.A. The isolated flat silicon nanocrystals (2D

structures) stabilized with perfluorophenyl ligands, Journal

of Nanoparticle Research 16(1): 2190, 2014 DOI: 10.1007/s11051‐

013‐2190‐4

IV Tonkikh A. A., Rybkovskiy D. V., Orekhov A. S., Chernov A.

I., Khomich A. A., Ewels C.P., Kauppinen E. I., Rochal S. B.,

Chuvilin A.L., Obraztsova E. D. Optical properties and

charge transfer effects in single‐walled carbon nanotubes

filled with functionalized adamantane molecules, Carbon 109:

87‐97, 2016 DOI: 10.1016/j.carbon.2016.07.053

V Kleshch Victor I., Tonkikh Alexander A., Malykhin Sergey

A., Redekop Eugene V., Orekhov Andrey. S., Chuvilin

Andrey L., Obraztsova Elena D., Obraztsov Alexander N.

Field emission from single‐walled carbon nanotubes filled

with CuCl, Applied Physics Letters 109: 143112, 2016 DOI:

10.1063/1.4964273

VI Bokova‐Sirosh S. N., Pershina A. V., Kuznetsov V. L.,

Ishchenko A. V., Moseenkov S. I., Orekhov A. S.,

Obraztsova E. D. Raman Spectra for Characterization of

Onion‐Like Carbon, Journal of Nanoelectronics and

Optoelectronics 8: 106–109, 2013 DOI: 10.1166/jno.2013.1444

VII Alexeev Andrey M., Ismagilov Rinat R., Ashkinazi Evgeniy

E., Orekhov Andrey S., Malykhin Sergei A., Obraztsov

Alexander N. Diamond platelets produced by chemical

vapor deposition, Diamond & Related Materials 65: 13–16,

2016 DOI: 10.1016/j.diamond.2015.12.019

VIII Orekhov Andrey S., Tuyakova Feruza T., Obraztsova

Ekaterina A., Loginov Artem B., Chuvilin Andrey L.,

Obraztsov Alexander N. Structural defects in single crystal

diamond needles, Nanotechnology 27(45): 455707, 2016 DOI:

10.1088/0957‐4484/27/45/455707

The above publications have been included at the end of this

thesis with their copyright holders’ permission.

AUTHOR’S CONTRIBUTION Author has planned and performed the experimental parts of

transmission electron microscopy study, analyzed and

interpreted data received, contributed in discussions and in

writing the experimental part of the papers I, II, IV, V, VI, VII,

VIII. The main ideas were generated in fruitful discussions of all

co‐authors team. The papers III and VIII were written by the

author. The experimental setup discussed in these papers,

samples preparation was performed by the author.

Interpretation of results which most important in electron

microscopy experiments were done by author on basic of

proposed atomic models materials studied. For these model

series of HRTEM images and electron diffraction patterns were

calculated.

Contents

INTRODUCTION ............................................................................. 1

1 Review of the structural features of carbon materials ............. 3

1.1 Types of hybridization of atomic orbitals in carbon ................ 3

1.2 Graphene and its derivatives ...................................................... 7

1.3 Methods of synthesis .................................................................. 20

1.4 Techniques for carbon materials structural characterization27

Objectives of this work ..................................................................... 34

2 GENERAL PRINCIPLES AND PRACTICAL ASPECTS OF

TEM ANALYSIS .............................................................................. 35

2.1 Theoretical aspects of electron microscopy ......................... 35

2.2 Optimization of spherical aberration correction for nano‐

carbon materials (Experimental part) ............................................ 55

Summary of the Chapter 2 .............................................................. 66

3 ATOMIC ARRANGEMENT OF GUEST CRYSTAL IN THE

INNER CHANNEL OF FUNCTIONALIZED SINGLE‐

WALLED CARBON NANOTUBES ............................................. 67

3.1 Singe walled carbon nanotube functionalized by adamantane

molecules ........................................................................................... 68

3.2 Structure characterization of Hg2Cl2 crystals located in the

inner channel of SWCNT ................................................................. 70

3.3 Crystal structure of 1D CuCl@SWCNTs .................................. 73


4 THE STRUCTURAL PECULIARITIES OF NANO‐ AND

MICRO‐METER SIZE SCALE DIAMOND CRYSTALS ......... 79

4.1 The transformation of Sp3 to Sp2 C‐C bond in nanodiamond

under heat treatment ........................................................................ 80

4.2 Diamond platelets produced by chemical vapor deposition 82

4.3 Structural peculiarities of single crystal diamond needles of

nanometer thickness ......................................................................... 84


5 CONCLUSIONS ........................................................................... 89

REFERENCES ................................................................................... 91

ORIGINAL PUBLICATIONS ..................................................... 105

Dissertations in Forestry and Natural Sciences Number 258 1

Introduction

Carbon materials always attracted great attention of

research community because of their importance for practical

applications and for fundamental science. During last decades,

especial interest has been dedicated to production and

comprehensive characterization of nano‐structured forms of

carbon materials such as nanodiamonds, fullerenes, carbon

nanotubes, graphene, numerous graphene‐like species and their

composites. These relatively “new” types of carbon materials

exhibit numerous unique properties promising important and

breakthrough achievements in fundamental science and

technology. Study of nanocarbon materials structural

peculiarities is necessary for deeper understanding of the

mechanisms determining their physical properties and for

optimization these properties for practical application purposes.

In spite of intensive study in this area many problems remain to

be open especially for specific forms of the nanostructured

carbon materials. This work presents results of high resolution

transmission electron microscopy analysis of structural

peculiarities of some of nanocarbons materials forms.

We performed optimization of optical system of

aberration corrected transmission electron microscope (TEM).

This adjustment allowed to increase of the signal‐to‐noise ratio

up to 30 % for the graphene TEM images with sub‐angstrom

spatial resolution at 80 kV accelerated voltage.

The optimized TEM instruments were used in this work

for structural characterization of the nanodiamonds, two‐

dimensional (2D) structures and needle‐like diamonds, onion‐

like nanocarbons and other graphene‐based structures,

including composites consisting of linear (1D) CuCl and Hg2Cl2

crystals encapsulated in single‐walled carbon nanotubes.

Obtained results allowed appropriate development of

Andrey Orekhov: Electron Microscopy Study of Structural Peculiarities of

Carbon Materials

2 Dissertations in Forestry and Natural Sciences Number 258

production processes for these carbon nanostructures and


In this Thesis after brief Introduction we are reviewing in

Chapter 1 most common structural features and methods used

for synthesis of carbon nanomaterials. Thereafter in Chapter 2

we are considering methodology of TEM experiments with

special emphasis to TEM instrument adjustments which are

necessary to achieve suitable reproducibility and interpretation

of obtained results. Chapters 3 and 4 present description of

obtained original experimental results and conclusions made on

their basis which are presented together with copies of the

articles where they were published.


1 Review of the structural

features of carbon

materials

Carbon materials have a wide range of properties which

make them suitable for various applications. Depending on type

of interatomic bonding and on crystal structure, carbon

materials could possess very different mechanical, thermo‐

physical, optical and electronic properties. Thus, diamond is the

hardest material which is used for indenters, anvils and similar

applications. On the other hand, graphite softness allows to use

it in gaskets and in lubricants. Diamond is transparent dielectric

material with the highest thermal conductivity. Conversely,

graphite is opaque, has good electrical conductivity and may be

used (in some cases) as thermal insulator [1], [2]. The

information about crystal structure composition and presence of

lattice defects is especially essential for carbon materials for

development of their production methods and for their

applications.

The following paragraphs present brief overview of the

structural features of carbon materials, the methods of their

synthesis and the most common techniques used for structural

characterization.

1.1 TYPES OF HYBRIDIZATION OF ATOMIC ORBITALS IN CARBON

The most common model describing interatomic bonding

assumes, so called, hybridization of atomic orbitals which

allows explanation of electronic structure formation of


Carbon Materials


molecules. In particular, it explains modification of atomic

orbitals during formation of a covalent chemical bond,

alignment of the lengths of chemical bonds and bond angles in

the molecule. Linus Pauling has proposed the scheme of valent

atomic orbitals hybridization. Spatial orientation of the hybrid

orbitals is defined the type Each type of hybridization. These

hybridized orbitals model is used for carbon material

characterization. Four valence electrons of carbon atom in its

ground state are arranged in one ‘2s’ orbital and two ‘2p’ orbitals.

With excitation of the atom one of two ‘2s’ electrons moves to

the free ‘2p’ orbital. The process of hybridization can be

represented as different combinations of one ‘s’ and one, two or

three ‘p’ orbitals. Thus there is potential possibility to form two,

three or four new orbitals. Each of these formed orbitals retains

some part of properties of the ‘s’ orbital and some part of

properties of the ‘p’‐orbitals. These new orbitals are called ‘sp1’,

‘sp2’, ‘sp3’ hybridized orbitals, correspondingly [3], [4], [5].

Orientation of the hybrid orbitals is determined by Coulomb

interaction between electrons and condition of minimization of

free energy of the system:

sp1‐hybridization occurs when one s and one p orbitals are

mixed; these two equivalent sp‐atomic orbitals have axial

symmetry and are arranged linearly with an angle between the

axis of 180 degrees and with opposite directions from the core of

the central atom; the remaining two non‐hybrid p‐orbitals are

perpendicular to the hybrid one and to each other; these non‐

hybrid orbitals provide formation of π‐bonds (Fig. 1‐1a);

sp2‐hybridization occurs when one of s‐ and two of p‐orbitals

are mixed; the obtained hybrid orbitals are arranged in one

plane with their axes directed from the atom to the vertices of

the triangle at an angle of 120 degrees; the non‐hybrid p‐orbital

is situated perpendicular to the plane and is may be involved in

the formation of π‐bonds (Fig. 1‐1b);

sp3‐hybridization occurs when one of s and all three p

orbitals are mixed to form four equivalent in form and energy

hybrid orbitals; the axes of the sp3‐hybrid orbitals are directed

from the atom toward the corners of a tetrahedron; the angle

Chapter 1: Review of the structural features of carbon materials


between any two axes is approximately 109°28ʹ, which

corresponds to the lowest electron repulsion energy (Fig. 1‐1c).

Figure 1‐1. The scheme of hybridization of atomic orbitals corresponds to sp1‐ (a) sp2‐

(b) and sp3‐ hybridization (c). Hybridized orbitals are shown by red and non‐hybrid

orbitals are shown by blue.

Type of hybridization determines atomic arrangements in

condensed carbon materials which may be represented by

different variants and combinations of amorphous carbon, one‐

dimensional linear atomic structures (carbyne), planar two‐

dimensional structures (graphene) or bulky three‐dimensional

structures (diamond):

amorphous carbon consists of a mixture of individual atoms

and/or atomic clusters with different hybridizations; typical

examples of amorphous carbon are wood charcoal, coke, soot

and similar objects which frequently contain small fragments of

stacked planar structures with interplanar distance of 0,350‐

0,365 nm; at high temperatures amorphous carbon may be

transformed into crystalline phase of graphite with different

levels of ordering; consideration of the structural organization of

amorphous carbon and its ordering is out of the scope of this

work;

carbyne is carbon materials with sp1‐hybridization of atomic

orbitals [6],[7]; this type of hybridization leads to formation of

the polymer‐like chains; there are two possibilities for carbyne

formation ‐ in β‐carbyne carbon atoms are arranged with double


Carbon Materials


atomic bonds (=С=С=С=) and in α‐carbyne with alternating of

single and triple bonds (‐C≡C‐C≡C‐); the polymer chains have

chemically active ends with localized negative charge; the

chains may be arranged by different manners to provide further

structurization; however this type of carbon materials is not

considered in this work also;

graphene is a monolayer of carbon atoms bonded together by

sp2 hybrid orbitals; orientation of the atomic orbitals leads to

formation hexagonal (or honey‐comb) atomic structure of

graphene; the non‐hybridized p orbital of each atom is oriented

perpendicular to the atomic plane and participate in formation

additional interatomic π‐bonding; similar to carbyne all atomic

bonds in graphene are saturated except dangling bonds on its

periphery; little variation of the angles between axes of the

hybridized orbitals make possible formation other numerous

types of carbon materials with spherical, cylindrical, conical,

wavy shape of the atomic layer; chemical interaction, which is

possible for atoms located in planar part of the atomic layer,

results in formation 2D corrugated structure of graphene; weak

Wan der Waals interaction between the layers is responsible for

coupling of graphenes into multilayer graphite or into

multishell spherical and cylindrical structures; some of these

graphenic structures will be analyzed in this work;

diamond structure is formed by carbon atoms with sp3

hybridized atomic orbitals; tetrahedral orientation of these

orbitals is responsible for 3D atomic arrangement in diamond

structure; at the same time each atomic layer in diamond

structure is similar to graphene sheet, which is corrugated due

to bonding with other carbon atoms, and, thus, ʹdiamondʹ

structure may be considered as stacked corrugated graphenes;

in this work we will consider some examples of diamond

structures.



1.2 GRAPHENE AND ITS DERIVATIVES

Graphene and its numerous derivatives attract especially

great attention after pioneering works of Andrey Geim and

Konstantin Novoselov from Manchester University. In 2004 they

extracted a single sheet of graphene using a micromechanical

cleavage technique [8]. Later in 2005, the same group measured

the quantum Hall effect and experimentally proved linearity of

energy spectrum of electrons in graphene and applicability of

Dirac equation for the charge carriers in graphene [9]. However,

graphene has been known long before these works. In form of

graphene oxide it was first discovered in 1859 by chemist

Benjamin Brodie. Late G. Ruess and F. Vogt have detected

evidence of presence of carbon crystals of atomically small

thickness by Transmission Electron Microscopy. Although, these

crystals have been not pure graphene and their thickness have

been of several nanometers. The first graphene layers on metal

(Ru, Rb, Ni) substrates have been obtained in 1970s by Blakely et.

al.[10], [11]. The history of graphene is started much before 2004

also because it has been considered for a long time as a model

for structural analysis of carbon allotropes including fullerenes,

carbon nanotubes and graphite. In these studies many

properties of graphene such, as a thinnest, strongest and perfect

2D crystalline structure with very high electron mobility and

thus it excellent conductivity of both heat and electricity, have

been analyzed. And at present days graphene become an ideal

object for both fundamental studies and for variety of exciting

applications.

Let’s consider crystal structure of graphene which essentially

determines all its properties. Crystal lattice of graphene,

schematically shown in Fig. 1‐2a, consists of regular hexagons.

The unit cell of graphene could be represented as two sublattices

of red (A) and blue (B) carbon atoms. Each equivalent atoms

participates in formation of a triangular sublattice which may be

described by translating vector , = m + n , where m and

n – integers, and and are Bravais crystal vectors for a unit

cell (marked as yellow rhombus). The vectors , and


Carbon Materials


specify position of atom B around site of atom A. Distance

between the nearest carbon atoms in the hexagons, marked a, is

0.142 nm. The unit cell parameter can be derived by ao=√3 , it is

equal to 0.246 nm.

Red tetravalent carbon atom A is covalently linked to three

neighboring blue carbon atoms B arranged in a plane, so the

angle between the bonds is 120°. The fourth electron occupies

2pz orbital, which is oriented perpendicular to the plane of the

graphene. These electrons are responsible for the unique

electronic properties of graphene and form a π‐zone in the

electron energy spectrum. Due to weak interaction of the π‐

electrons with a particular atom they may be easily moved

along graphene plane providing its high electrical and thermal

conductivity.

Figure 1‐2. Schematic representation of the hexagonal structure of graphene with unit

cell marked by yellow rhombus (a). Atomic model of perfectly flat sheet of graphene (b)

and the wavy roughness (c) at the molecular level found experimentally (adopted from

[12] with permission num. 3997280548454).

The atomic structure of an isolated single‐layer graphene has

been studied by transmission electron microscopy (TEM) [12].

Obtained electron diffraction patterns evidenced about expected

honeycomb crystal lattice. Detailed analysis of the diffraction

spots intensities reveals that the graphene sheet has a wavy

roughness with amplitude of about one nanometer (Fig. 1‐2c).



This roughness may be due to instability of the two‐dimensional

crystals [13–15]. It is believed that only through this wavy

roughness the graphene sheet could exists in a free‐standing

form and do not curl up spontaneously.

Despite of detailed analysis in previous studies many

problems related to identification and explanation of structural

peculiarities of graphenic materials require further research. The

interest to this research is stimulated by recent technology

developments which make possible reproducible production of

nanostructured materials with unique properties and their

practical usage. This includes purely carbon materials (such as

graphene and graphite, carbon nanotubes and fullerenes,

diamonds) as well as their different composites (like

heterostructures, filled single‐ and multi‐walled carbon

nanotubes).

1.2.1 Graphite

The most typical example of graphene‐based materials is

graphite. The structure of graphite consists of stacked parallel

graphene layers. Because of saturation of all valence bonds of

carbon atoms in graphene, the parallel atomic layers in graphite

are joined together only by weak Wan der Waals interaction. It

provides one of the most distinguishable properties of graphite

which consists in easily detaching material from a piece of

graphite when it is moving in contact with a surface. The

stacked atomic layers have some planar displacement from each

other and probability to form certain type 3D structure is

determined by free energy of the system. The most common (i.e.

the most probable) forms of graphite have hexagonal and

rhombohedral crystal lattice. The hexagonal lattice is formed by

alternation of A and B graphene layers in ‐A‐B‐A‐B‐ stacking

(Fig. 1‐3). A and B atoms on consecutive layers are on top of one

another (marked by yellow color) and A’ atoms in one plane are

over the hexagon centers of the adjacent layers and similarly for

the B’ atoms [16]. This graphene layer arrangement is called

Bernal stacking. In‐plane the nearest neighbor distance ac‐c of

0.142 nm, in‐plane lattice constant ao of 0.246 nm, and c‐axis


Carbon Materials


lattice constant co of 0.670 nm. This structure belongs to P 63/mmc

space group and has four atoms per unit cell (see Fig. 1‐3b).

In the rhombohedral (β‐phase) graphite the alternation

follows to the ‐A‐B‐C‐A‐B‐C‐ stacking order. It has ao of 0.246

nm, co of 1.005 nm and belongs to R‐3m space group with six

atoms per unit cell. This structure of β‐graphite is metastable

unlike to more stable α‐phase of graphite. In graphite of natural

origination about 30% of content may have rhombohedral

structure. Artificially synthesized graphite contains practically

only hexagonal form.

Figure 1‐3. Scheme of two consecutive graphene layers arrangement in hexagonal

graphite structure. A and B atoms are on top of one another (depicts by yellow color)

and A’ atoms in one plane are over the hexagon centers of the adjacent layers (a). The

hexagonal unit cell selection is represented in (b).

The crystal structure of hexagonal graphite belongs to

P63/mmc space group [17]. The symmetry elements of this space

group are schematically represented in Fig. 1‐4 [18]. This space

group includes following axes: vertical sixfold screw axis 62

which is a composition of a rotation by an angle 60° followed by

a translation along this axis of 1/3 of the lattice vector c,

inversion threefold axis, vertical twofold screw axis 21 which is a

composition of a rotation by an angle 180° followed by a

translation along this axis of ½ of lattice vector c, horizontal

twofold axis 2 and horizontal twofold screw axis 21 with the

translation along the horizontal coordinate and diagonal

directions. The group P63/mmc also contains symmetry planes:

horizontal mirror plane, vertical mirror planes which alternate



with glide planes b (composition of the reflection and translation

along half the lattice vector b) and vertical glide planes c which

alternate with n glide planes (composition of the reflection and

translation along half a face diagonal). The interaction of the

twofold axes and planes perpendicular to them leads to the

appearance of the centers of symmetry. Carbon atoms occupy

two types of symmetrical non‐equivalent positions which are

characterized by Wyckoff positions 2b and 2c. These initial

atoms are multiplied by all symmetry elements of the space

group in a way that carbon atoms are arranged in hexagonal

layers alternating along lattice vector c in a checkerboard

pattern.

Figure 1‐4. Graphical representation of symmetry elements of crystal structure with P

63/mmc space group (adopted from [18]. Reproduced with permission of the

International Union of Crystallography).

Table 1‐1. Crystal unit cell parameters for known graphite phases.

Modification a, b, c, (nm)

α, β, γ Space Group

Cryst. Sys. ICSD ref.

α-phase (2H) 0.246 0.246 0.670

90.00 90.00 120.00

P63/mmc hexagonal 52230

β-phase (3R) 0.370 0.370 0.370

39.75 39.75 39.75

R3mR rhombohedral 53780

The crystal lattice symmetry determines physical properties

of graphite. Particularly the properties of graphite have strong


Carbon Materials


anisotropy which is determined by its layered structure [19],

[20]. Also presence of symmetry center leads to absence in

graphite polar directions and, consequently, physical properties

which may be assigned to the polar symmetry characteristics.

The unit cell parameters, space group and references on

Inorganic Crystal Structure Database (ICSD) for hexagonal α‐

phase (2H) and rhombohedral β‐phase (3R) graphite phases are

presented in Table 1‐1.

1.2.2 Single and Multi‐Walled Carbon Nanotubes

The nanotubes are graphene layers rolled into a cylinder‐type

structure with diameter of several nanometers. Rolling of a

single graphene sheet leads to formation of single‐walled carbon

nanotubes (SWCNT), whereas multi‐walled carbon nanotubes

(MWCNT) consist of several coaxial single‐walled nanotubes.

Nanotube diameter corresponds to geometry of the graphene

honey‐comb lattice and twisting angle of the rolled graphene

sheet. Twisting type is specified by the chiral vector [21], [22],

[23].

The diameter and chirality of the nanotube can be expressed

using by vector Ch = n + m ≡ (n, m). This vector connects

two crystallographically equivalent sites in the 2D graphene.

The and are the graphene lattice vectors, n and m

determine the diameter and the twisting angle of the nanotube.

If m = 0, carbon nanotubes belong to the zigzag type if n = m, the

nanotubes belong to armchair type. Otherwise, when n ≠ m they

are called as chiral tubes (Fig 1‐5). Armchair and zig‐zag type

tubes structure have a high symmetry. These terms refer to

location of the hexagons along the circumference. While chiral

type of tube means, that tube can exist in two mirror‐related

forms.

Diameter of carbon nanotubes can be calculated from n, m as

follows:

Where a is length of unit cell vector a1 or a2. Length of a is related to carbon‐carbon bond length: a = |a1|=|a2|=acc√3. For graphite carbon‐carbon bond length is acc = 0.142 nm. But for



nanotubes due to curvature of atomic sheet this value slightly

large as acc = 0.144 nm [24]. Nanotube sizes may vary in a wide range. The minimal

diameter of SWCNTs is 0.4 nm and maximal diameter does not

exceed 2‐3 nm. Length of nanotubes may reach up to several

mm. Outer diameter of multiwalled nanotubes may be from 4 to

30 nm and its length may be reach several μm. Inter‐walled

distance in MWCNT is similar to inter‐planar spacing in

graphite and equals to 0.335 nm.

Figure 1‐5. Schematic illustration of carbon nanotube (SWCNT and MWCNT)

formation by rolling up graphene sheet with different chiral vector: n=m – armchair

tubes, m=0 zigzag tubes and n≠m chiral tubes.

The values of n and m determine the chirality of nanotube

and affects to the electronic, optical and mechanical properties.

The differences of interatomic bonds angles in the planar atomic

structure of graphene and in the nanotubes lead to modification

of electronic properties. These differences are negligible for

direction along the nanotube axis. So, quasi‐one‐dimensional

character of structural and electronic properties of the

nanotubes is observed. Also variation of twisting angle and

diameter of the nanotubes produce their different electronic

properties. In particular, the twisting determines conductivity

type and band gap of the nanotubes. Nanotubes with |n—m| =


Carbon Materials


3i (i is an integer) possess a metallic properties while for |n—m|

= 3i ± 1 the nanotubes are semiconductors. Recent advances in

nanotube study are discussed, for example, in papers [25], [26].

1.2.3 Fullerenes

Defects of atom packing in planar graphene structure may

lead to distortion of an ideal honeycomb atomic arrangement

and formation of pentagon and heptagon‐like configuration of

carbon bonds [27], [28]. The specific sequence of pentagon and

heptagon arrangement results in keeping of sheet planar

structure, but may initiate grain boundaries formation. The

individual grains may be stitched together in graphene sheet

through the pentagon‐heptagon pairs. In other case

geometrically noncompensated defects may cause formation of

roughness on graphene sheet as well as formation of spherical

structures. Typical example of such kind of structural

modification is formation of closed structure of fullerene.

Similar to carbon nanotubes, possible sizes and geometries of

the fullerenes are completely determined by geometry of carbon

net in graphene. Stability of the fullerene molecules is

determined by a number of distorted interatomic bonds and

thus decreases with increase of their atomic weight (i.e. increase

of number of carbon atoms or diameter of the molecule). This

tendency determines existence of optimal diameter of the

fullerene which corresponds to the most stable and the most

probable configuration. This diameter equals about 0.7 nm.

Number of carbon atoms which is necessary to form such

fullerene is 60 and they should be arranged into 12 pentagons

and 20 hexagons [29]. This is the most common type of fullerene

– C60 (see Fig. 1‐6). Other fullerenes may be produced and

obtained as pure fraction also but their stability and yield in

synthesis process will be less in comparison with C60. The

smallest available fullerene consists of 20 carbon atoms and its

diameter is about 0.3 nm. Larger fullerenes are also possible (the

most common fullerenes contain up to 100 carbon atoms).

Similar to multi‐walled carbon nanotubes and to multilayer



graphite structure there are fullerenes consisting of few co‐

centric shells called also as onion‐like molecules.

The deviations of interatomic bonding which happens in

fullerenes in comparison with that in planar graphene, induce

formation of unique electronic and chemical properties in these

spherical structures, which are attractive for numerous

applications. Similar modifications are possible also with

introduction into hexagonal network of undisturbed graphene

not only pentagons but other polygons (e.g. heptagons) also [30].

Figure 1‐6. Schematic presentation of defects in the ideal graphene package – pentagon

(a) and heptagon (b) atomic arrangements. C60 molecular structure consisting of 12

pentagons surrounded by 20 hexagons (с).

An example of packing defects in graphene sheet is shown in

Fig. 1‐7. The series of high resolution TEM images represent the

transformation of C‐C bonding appeared under electron beam

irradiation. The pentagons and heptagons formations are

occurred in the honey comb structure. The defects formation

mechanism in graphene sheets is reviewed in detail in [12,31,32].

Figure 1‐7. A series of high resolution transmission electron images of graphene sheet

with packing defect.


Carbon Materials


1.2.4 Functionalized Graphene

Alternatively, the distortions in interatomic bonds (and

corresponding properties modifications) of flat graphene may

be achieved via interactions of carbon atoms with other atoms or

molecules. This modification of graphene is called as its

functionalization [33], [34], [35]. There are two main categories

of functionalization: chemical and nonchemical. Chemical

functionalization is realized through the formation of new

covalent bonds between carbon atoms and the guest functional

groups; nonchemical functionalization is mainly based on π

interaction between unhybridized orbitals of carbon atoms and

guest molecules i.e., mainly a physical interaction. For example,

a regularly corrugated structure of the graphene sheet may be

formed, as a result of chemical reaction with suitable precursors

when three hybridized bonds of carbon atom are shifted from

graphene plane. This type of structural modification opens

attractive possibility for creation the heterostructures consisting

of the stacked functionalized graphene sheets. An example of

graphene structure modification due to its functionalization is

shown in Fig. 1‐8. The image shows schematic model of

hydrogenated graphene (also called as graphane). Advances in

the graphene functionalization are well described in the review

[36].

Figure 1‐8. Graphane atomic structure as an example of graphene structure

modification due to its functionalization.



1.2.5 Diamond

Atomic arrangement in individual atomic layers of diamond

is similar to graphane (i.e. hydrogenated graphene). This is

because of sp3 hybridization which is necessary to provide

covalent bonding with participation of forth valence electron.

Thus 3D atomic structure of diamond may be represented as

stacked corrugated graphane layers where hydrogen atoms are

replaced by carbons from adjacent layer (Fig. 1‐9). This

approach naturally explains, for example, difference in

electronic properties of diamond and graphite by

transformations which happened due to deviation of the

interatomic bonds from planar graphene configuration. Crystal

structure of diamond, resulted of this, is described by a face‐

centered cubic (or FCC) lattice.

Figure 1‐9. Schematic representation of cubic diamond crystal structure by ‐A‐B‐C‐

stacked corrugated graphene layers. Alternatively diamond structure may be

represented by combination of tetrahedrons.


Carbon Materials


There are four types of the diamond phases: cubic phase 3C

[37], [38] and tree hexagonal modifications 4H, 6H, 8H

(lonsdaleite) [39], [40]. Crystallographic parameters of these

phases are represented in Table 1‐2. Cubic diamond structure

can be considered as corrugated graphene layers arranged in –

A‐B‐C‐ sequence normal to [111] direction. Otherwise, structure

of diamond can be represented as a combination of two

interpenetrating FCC sub lattices which are displaced along the

body diagonal of the cubic cell on 1/4th length of the diagonal.

Each carbon atom is joined with four other carbon atoms

forming a regular tetrahedron with the distance of 0.154 nm,

angle between the carbon‐carbon bonds is 109.28o.

Cubic diamond belongs to 3 space group. The symmetry

elements of this space group are schematically represented in

Fig. 1‐10 [18]. This space group include following axis: fourfold

screw axis with 1/4 shifting along c, inversion fourfold axis

perpendicular to a mirror plane, treefold rotation axis, threefold

screw axis with 1/3 shifting along c, twofold rotation axis,

twofold screw axis with 1/2 shifting along c. 3 consists also

of planes of symmetry: glide planes along and perpendicular to

c axis, mirror planes along to c axis and diagonal glide planes.

Figure 1‐10. Graphical representation of symmetry elements of crystal structure with

F d 3 m space group (adopted from [18]. Reproduced with permission of the

International Union of Crystallography).



The crystal lattice symmetry and strong covalent interatomic

bonding determine physical properties of diamond. For

example, diamond is the hardest known naturally occurring

material. Diamond crystals predominantly are good electrical

insulators. But diamonds with substitutional boron impurities

reveal semiconductor properties. Diamond also is a good

conductor of heat because of strong covalent bonding and low

phonon scattering.

At the same time, interatomic bonding deviations from flat

graphene configuration leads to increase of specific free energy.

Thus, diamond structure possesses metastable characteristics.

However, under normal conditions diamond will not decay into

graphite due to large potential barrier for this transition.

Table 1‐2. Crystal unit cell parameters for known diamond phases.

Modification a, b, c (nm) α, β, γ Space

Group Cryst Sys ICSD ref.

Diamond 3C 0.356 0.356 0.356

90.00 90.00 90.00

F d 3 m cubic 44100

Diamond 4H 0.252 0.252 0.823

90.00 90.00 120.00


Diamond 6H 0.252 0.252 1.235

90.00 90.00 120.00


Diamond 8H 0.252 0.252 1.647

90.00 90.00 120.00


1.2.6 Filled Carbon Nanotube

Closed structure of fullerenes and carbon nanotubes allows

alternative possibilities for variation of their properties (first of

all electronic) via filling by different components [41], [42]. In

contrast to that considered above, it may be that there are no

new covalent bonding between carbon atoms and filling

precursors [43]. However, modification of electronic properties

and structural parameters might be quite large. It is possible to

achieve the structural ordering of the filled compound inside the

fullerenes or nanotubes. During metallic nanotubes filling by

electron‐donor with Fermi level higher than SWCNT level the


Carbon Materials


electron density of the nanocomposite will increase. On other

hand, the nanocomposite will possess the semiconductor

properties when nanotube is filled by electron acceptor [44], [45].

The nanotube filling provides realization of attractive

possibility for chemical synthesis. In this case inner volume of

one‐dimensional carbon tubes is considered as a nanoreactor

[46], [47]. With filling this reactor by chemical reagents one can

obtain conditions for production, for example, new compounds

with untypical crystal structure and stoichiometry due to

encapsulated crystal‐nanotube interaction. It was shown that

structures of encapsulated crystal may differ from bulk crystal

structure with changing the crystal symmetry, bond lengths,

and bond angles. Geometrical distortions are caused by limit

inner nanotube space and adjustment of encapsulated crystal

structure to the internal diameter of SWNT channel.

1.3 METHODS OF SYNTHESIS

Structural and physical properties of nanocarbon materials

depend essentially on conditions of their synthesis. The most

general approach to carbon materials production assumes

condensation of carbon atoms (or clusters) from a vapor phase.

Thus, from this ’general’ point of view, the most important

synthesis process steps include:

(1) Formation of atomized (vaporized) carbon environment.

(2) Procuring sufficient activation energy to carbon atoms.

(3) Formation of appropriate type of carbon materials by

applying controlled condensation conditions.

Phase diagram of carbon (see Fig. 1‐11) predicts possibility

for obtaining diamond (sp3) type of carbons by transformation

of other condensed carbon containing materials (usually

graphitic – sp2) at certain conditions. Below we will consider

some typical growth methods which seem to be the most

representative in view of results discussed in this work.



Figure 1‐11. Phase diagram of carbon. The transformations between different phases

are possible with appropriate conditions.

1.3.1 High pressure, high temperature method

Method using high pressure and high temperature (HPHT) is

still widely used because of its relatively low cost. HPHT is a

relevant process for the synthesis of sp3 carbon structures as

well as other ultra‐hard materials as cubic BN with sp3

hybridization. In the HPHT method for obtaining the pressure

and temperature, which are necessary to produce synthetic

diamond, may use three press designs: the belt press, the cubic

press and the split‐sphere press [48], [49], [50]. The graphite

particles are generally used as the precursor which placed at the

bottom of the press. The internal part of press is heated above

1400 °C and melts the solvent metal. The molten metal dissolves

high purity carbon source. The dissolved carbon is then

transported towards to small diamond precursors and forms a

large synthetic diamond. Adding metallic solvents leads to an

important lowering of pressure and temperature required for

diamond synthesis. Without addition of those catalytic metals,

crystal structure of graphite remains unchanged after

compression up to 15 GPa at room temperature and loses some

of graphite features at higher pressure. Possible transformation

into metastable hexagonal diamond form (lonsdaleite) appears

to depend on many factors as temperature or degree of

disorganization of the initial sp2 carbon form.


Carbon Materials


1.3.2 Detonation of explosives

Methods of explosive synthesis of diamonds are similar to

HPHT with the high pressures and temperatures achieved

during short‐term detonation and subsequent rapid cooling of

carbon‐containing explosive material. Immediately with

diamond other carbon phases are formed in this process usually

because of wide spread of achievable pressure and temperature

parameters [51].

Pure graphite may be transformed into diamond by a shock

compression [52], [53], [54]. By this method transformation of

rhombohedral graphite into diamond by shock wave pressure at

30 GPa has been obtained. Exposure time has been a few

microseconds and temperature of 1300‐1800 K. Small diamonds

(2‐5 nm) have been produced. In another method diamond

crystals have been produced from graphite powder enclosed in

a metallic matrix undergoing by high pressures (9‐20 GPa) and

temperatures (2300‐3300 K) during microseconds. Cubic

diamond phase and up to 30‐50% hexagonal diamond phase

(lonsdaleite) have been formed by this techniques.

The short duration pressure and temperature increase up to

values corresponding to graphite‐diamond transformation may

be achieved also in synthesis processes based on usage of high

intensity pulse laser irradiation initiating explosion of graphite

dispersion in liquid xenon.

1.3.3 Arc‐discharge

There are numerous version of methods providing synthesis

of condensed carbons from a vapor phase. One of the most

simple (and thus popular) is the arch‐discharge methods [55],

[56]. In accord with this method synthesis occurs in a chamber

with two electrodes, one of which contain a powdered carbon

precursor. The pre‐evacuated chamber is filled with inert gas

(like helium or argon) at low pressure (between 50 and 700

mbar) (Fig. 1‐12). One of the electrode ‐ anode may be arranged

axially and moveable relative to other ‐ cathode in longitudinal

direction. To initiate arc discharge the electrodes are moved

closer (typical distance is about 1 mm ). A constant current is



maintained through the electrodes

to obtain a non‐fluctuating arc. The

fluctuating arc results in unstable

plasma and affects on quality of

the synthesized product. The arc

discharge current generates

plasma of very high temperature

~4000–6000 K, which sublimes the

carbon precursor, which are filling

the anode. The vaporized carbon

atoms (clusters) aggregate in the

gas phase and drift towards the

cathode where they are cooled down due to the temperature

gradient. During the condensation different types of carbon

clusters, including fullerenes and/or carbon nanotubes, may be

obtained depending on process conditions. However nanotube

production requires usage of catalysts particles which may be

produced immediately in the gas discharge plasma [57][58]. A

comprehensive review of arc synthesis technique is done in [59].

1.3.4 Laser Ablation

There are many other variants

of techniques providing

vaporization and excitation of

carbon atoms via different types

of gas discharge or by thermal

evaporation. For example, in the

laser ablation method intense

laser pulses are used to heat and

ablate a carbon target, which is placed in a tube‐furnace heated

to 1200°C (Fig. 1‐13). During the process inert gas like helium or

argon is flowed through the chamber to carry the grown

nanotubes to the copper collector. After cooling of the chamber

nanotubes and by‐products, like fullerenes and amorphous

carbon, can be collected [60], [61], [62].

With pure carbon laser ablation methods leads to synthesis of

multi‐walled nanotubes while addition of a catalyst like iron,

Figure 1‐12. A schematic

illustration of arc‐discharge

technique.

Figure 1‐13. A scheme of laser

ablation synthesis technique.


Carbon Materials


yttrium, sulphur, nickel and molybdenum leads to formation of

single‐walled carbon nanotubes [63], [64].

1.3.5 Chemical vapor deposition (CVD)

Essential peculiarity of chemical vapor deposition (CVD)

methods (from point of view of this particular work) consists in

its ability to produce thin film materials on suitable substrates.

In this method the film growth occurs under low pressure

(below atmospheric pressure). The method involves feeding a

mixture of gases (typically methane and hydrogen) into a

chamber and their decomposition onto chemically active

radicals. Potentially, for activation of the gas mixture different

methods can be used including microwave or direct current gas

discharges, optical discharge in focused laser beam, acetylene

torch or gas decomposition on ʹhot filamentʹ, etc. These methods

are mostly used for production of polycrystalline diamond films.

These diamond (i.e. mainly sp3 carbon) films usually contain

some amount of graphitic (i.e. sp2 carbon) and amorphous

carbon material. Depending on the process conditions the

percentage of different phases (sp3, sp2 and amorphous) in the

films may be changed in a wide range [65].

Figure 1‐14. Scheme of the CVD system with gas mixture activation by direct current

discharge.

There is great variety of technical solutions allowing

realization of the CVD process. Let us consider, as an example,

with more details, CVD system which has been used for

production materials studied in this work. Schematically this

system is presented in Fig. 1‐14.



The reactor chamber of the system is made of stainless steel

with double walls with water cooling system. Inner diameter

and height of the chamber is 90 cm. Two quartz windows in the

reactor chamber are used for controlling of the deposition

process including substrate temperature measurement by

optical pyrometer. The substrate is located on the sample holder

which is electrically connected with ground. Other electrode,

which is electrically connected to negative source of the power

supply, is located above the substrate on distance of about of 40

to 50 mm. The substrate holder (which plays role of anode, i.e.

positive electrode in electric circuit) and cathode (i.e. negative

electrode in the circuit) are made of molybdenum and have

shape of disks with diameter of 50 mm and thickness of 10 mm.

The disks are tightly attached to the holders cooled by flow

water.

After substrate is loaded the chamber is pumped out by

rotary pump up to vacuum level of about 10‐3 Torr. After that

the chamber is filled by a mixture of hydrogen and methane

with predetermined proportion depending on type of produced

material. Typical pressure of the gas mixture during deposition

is about 100 Torr. The direct current (DC) discharge is ignited

between the electrodes by applying voltage. The discharge

plasma contains chemically active radicals produced from

hydrogen and methane molecules. The system contains control

elements allowing support of necessary discharge voltage and

current, total gas mixture pressure and composition during long

time (up to hundreds hours) which is necessary for production

carbon films of different thickness [66].

The CVD method is widely used for the diamond fabrication

[67], [68], graphene synthesis [69] and production of nanotubes

[70], [71].

1.3.6 Methods of carbon nanotubes filling

There are two approaches to production of filled carbon

nanotubes. In the first approach the encapsulation of the

inorganic elements takes place during nanotube synthesis (in‐

situ methods). This method can be used to encapsulate metal


Carbon Materials


carbides or pure elements (Se, Ge, Sb, Cr, Mn, Co, Cu, Re, Au,

Sm, Gd, Dy, Yb) into nanotube under special synthesis

conditions in presence of catalyst. The disadvantage of this

approach is problem of control the filling process due to large

temperature gradient in the cathode area. Another disadvantage

of this method of synthesis is that it does not allow filling of

nanotubes with transition metals, as well as low efficiency of the

synthesis: the yield of filled nanotubes is only several percent

[43], [72], [73], [74].

Second approach is ex‐situ methods, when already

synthesized carbon nanotubes are undergo the filling. Special

treatment of the nanotubes is necessary to open their ends

(which are usually capped in as‐grown material). Depending on

type of desired compound to be filled into the nanotubes there

are few techniques are used for the filling.

Gas phase filling of CNT

Filling of carbon nanotubes from gas phase is carried out in

vacuum at high temperatures [75,76]. During nanotube heating

the low‐pressure vapor consisting of the compounds, which

should be encapsulated, undergoes capillary condensation and

thus penetrates into the nanotube. During subsequent cooling

the compounds are crystalized.

Liquid phase filling of CNT

Filling of SWNTs from liquid phase is performed using so‐

called capillary method, which involves impregnation of opened

nanotubes with solutions or suspensions of selected compounds

[47,77]. Higher concentration provides higher yield of this

method in comparison with usage gaseous environment.

Encapsulation from melts

This method uses liquid environment consisting of melted

components. The technique provides a 2‐3 times larger yield of

encapsulated nanotubes as compared to filling from

suspensions and solutions. The encapsulation procedure is



usually performed under vacuum conditions at temperatures

10‐100°C higher than the melting point of the guest material.

During subsequent slow cooling encapsulated material is

crystallized [43,78].

1.4 TECHNIQUES FOR CARBON MATERIALS STRUCTURAL CHARACTERIZATION

Structural and physical properties of carbon materials may be

characterized by common analytical methods used in material

science. Specificity some of these methods are considered below

briefly in view of their usage in this particular work. These

methods allow obtaining additional information for

identification of morphological features, crystal structure,

chemical composition, lattice defects, grain boundaries, phase

transitions, defects, and other parameters of studied materials.

This information can be obtained by applying a complex of

modern diagnostic methods including Raman spectroscopy (RS),

X‐ray diffraction (XRD), scanning electron microscopy (SEM),

energy‐dispersive X‐ray spectroscopy (EDS), electron diffraction

(ED), transmission electron microscopy (TEM), scanning

transmission electron microscopy (STEM) which are the most

widely applied techniques for characterization of carbon

materials.

1.4.1 Raman spectroscopy

Raman scattering is one of the most informative and widely

used non‐destructive methods for structural and chemical

diagnosis of various forms of carbon allotropes. Spatial

localization of phonons in nanostructured carbon materials

provides additional ability of RS to detect size dependent

structural changes and deformations in atomic bonding. Each

band in RS spectra is directly related to a specific frequency of

intramolecular vibrations. The vibration frequencies and, thus,

the positions of the Raman bands are very sensitive to


Carbon Materials


interatomic bonds orientation and to number of atoms at the

bond ends.

In Raman spectroscopy the sample is irradiated with

monochromatic light, produced by a laser source, and scattered

light is registered using high resolution spectrometers and

sensitive detectors [79], [80], [81], [82]. Practically, most of the

scattered light has the same frequency as the incident laser

radiation. Only very small amount of scattered light contains

information about studied matter, which is measured by a

frequency shift from the laser frequency. Fig. 1‐15 represents a

typical Raman spectrum of single‐walled carbon nanotube

which contain, so called, G (~1500 cm‐1), D (~1355 cm‐1) and

RBM‐bands. The G‐band is associated with optical vibrations of

two adjacent carbon atoms in lattice having ʺgraphiticʺ atomic

structure. D‐band indicates presence of structural defects. In

low‐frequency region Raman spectra could contain, so‐called

Radial Breathing Mode ‐ RBM‐band, which is a typical

characteristic of nanotubes. The RBM‐band is associated with

symmetric vibrations of carbon atoms in the radial direction.

The RBM‐band frequency is proportional to the nanotubes

diameter (~ 1/dNT). The G to D bands intensity ratio allow

estimation of density of structural defects. If intensity of D‐band

is significantly lower than intensity of G‐band this indicates

presence of sufficiently small amount of structural defects [76].

Figure 1‐15. An example of typical Raman spectrum of single‐walled carbon nanotube.

The structural models represent schematically atomic vibrations corresponding to

different RS modes.



Popularity of Raman spectroscopy in studies of carbon

materials is determined by important advantages of this method.

It requires little or no sample preparation and does not depend

of water absorption band. This allows carrying out

measurements of solids, liquids and gases through the

transparent containers such as glass, quartz and polymers.

1.4.2 X‐ray diffraction

X‐ray diffraction is a classical method for determining crystal

structure of different types of material. Structural studies of

carbon materials with usage of this method have been carried

out for a long time. This method allows distinguishing the

carbon allotropes, graphitic and amorphous carbon. When

carbon materials produce it, structure can vary over a wide

range depending on temperature and other conditions applied.

In particular, carbon atoms in plane of graphite layers can be

rotated or shifted laterally in respect of their ideal orientation in

the graphite structure. These detailed features are mirrored in

the X‐ray spectra. Estimation of diffraction peak width provide

information about coherent scattering region size and area with

large crystal distortions [83], [84].

The shape of diffraction peak can relay to phase composition.

So, the asymmetry of the diffraction peak profiles which

observed in some samples for both small and large scattering

angles is typical characteristic of multiphase materials. It may be

caused by presence of ordered and amorphous carbon in the

samples. X‐ray diffraction allows controlling functionalization

process of carbon materials. Due to functionalization some

characteristic peaks of spectra are shifted and changed their

parameters [85], [86].

The analysis based on observations of X‐ray diffraction has

some disadvantages. One of the most significant relates to size

of sample. While the method allows determination of sample

cell parameters with high accuracy volume of the sample should

be quite large that often difficult for nanomaterial studies.

Additionally, the method provides integral information, i.e. it

cannot be used for identification local structural distortions.


Carbon Materials


1.4.3 Scanning Electron Microscopy

Morphology, shape and size distribution of carbon‐contained

micro and nanoparticles, local and average microstructure

information, presence of pores, cracks and other defects – these

data are very important in structure characterization and

growth processes optimization of carbon materials [87], [88].

Scanning electron microscopy allows obtaining

microstructure information with high accuracy, as well as

integrated information from different areas of the sample. The

resolution of a SEM image depends on electron beam probe size

and on volume of interaction between the incident beam and the

sample. Microstructure information could be obtained from a

depth in range of tens of nanometers up to few microns

depending on e‐beam accelerating voltage (from 1 kV to 30 kV)

and on atomic number of sample material. The smallest electron

beam probe size and the highest signal to noise ratio is required

to achieve the SEM image with the best spatial resolution. The

smallest diameter of the probe size for the LaB6 cathode is ~ 5

nm while for field emission gun is ~ 5‐25 nm. The resolution of

the modern SEM instruments is about several Angstroms, which

is achieved with field emission electron source and special

design of optical system of electron column. SEM images could

be formed by registration of secondary and backscattered

electrons. Secondary electron imaging, being more surface

sensitive, has greater resolution. On the other hand,

backscattered electrons are sensitive to atomic mass of nuclei

they scatter from. Therefore, heavier elements which back

scatter electrons with high energy more efficiently appear

brighter than lighter elements. Thus SEM image in backscattered

electrons mode will possess a composite or Z‐contrast.

1.4.4 Energy‐Dispersive X‐ray Spectroscopy

The scanning electron microscopes (and transmission

electron microscopes (TEM) also) could be equipped with

facilities for chemical analysis of the samples by registration the

X‐ray signals generated in interaction of e‐beam with the matter.



EDS technique allows to carry out local qualitative and

quantitative chemical analysis [89], [90], [91].

The main advantage of X‐ray microanalysis of thin specimen

is improved spatial resolution over that obtainable in a bulk

sample. In the bulk samples normally analyzed in SEM, the

electron beam diffuses in the sample to a depth of 1‐5 μm. This

limits spatial resolution of EDS by about 1 μm for bulk samples.

In TEM thin foil the electron beam passes through the sample

and X‐rays are generated in a volume corresponded to the

focused probe size and the electron scattering area, which is a

function of both the thickness of the sample and the accelerating

voltage. The accelerating voltage for microanalysis in SEM is

choose to provide minimum volume of beam and sample

interaction which may be estimated by using Monte‐Carlo

simulation. On the other hand the energy should be sufficient to

excite the characteristic lines of the analyzed element.

The X‐ray qualitative analysis is based on the ability of a

spectrometer system to measure characteristic line energies of

specific elements. This analysis can now be performed with

accuracy about 1%, so measurement of the characteristic X‐ray

intensity should be not less than 1% in accuracy. The analysis includes Fourier filtering technique which is necessary to

remove background signal and to obtain intensities of the lines

corresponding to analyzed elements [92].

1.4.5 Transmission electron microscopy

Transmission electron microscopy is one of the most effective

methods of structural analysis. One of the greatest TEM

advantage is opportunity to analyze information both in real

and reciprocal space. TEM is widely used nowadays for analysis

different types of carbon nanocomposites and, in particular, for

filled nanotubes; for identification grain boundaries, twin

boundaries; for determination particle morphology, size

distribution, inclusions and other important structural

characteristics [93–95].

Typical range of accelerating voltages in a transmission

microscope is 80 ‐ 300 kV. It depends on sensitivity of the


Carbon Materials


analyzed sample to e‐beam irradiation. Typical resolution of

transmission microscope is about 0.2 nm for 300kV. The

resolution is reduced due to chromatic and spherical aberration

of the e‐beam optical system. Chromatic aberration contribution

can be decreased or minimized by applying electron source with

high coherency (typically field emission gun) or with use e‐

beam monochromator. More details related to this problem will

be discussed in Chapter 2. For minimization of spherical

aberration Cs‐corrector is used.

The procedures of the transmission electron microscopy

observation include wide range of techniques. In the bright field

TEM the image is formed by transmitted e‐beam. In this case

obtained contrast in TEM image is proportional to the thickness

and density variations in a sample. Dark field TEM allows

visualization of the particles or sample areas having different

compositions. Formation of the dark field image requires usage

of only single diffracted e‐beam. This techniques allows

detection of distribution of size and orientation of the particles

in the sample [96], [97]. Electron diffraction (ED) analysis can

provide important information about crystallinity of the

samples, i.e. about amorphous, polycrystalline and single crystal

structure. Phase identification of the samples can be revealed

through ED data, including interplanar distance, angles between

reciprocal vectors.

1.4.6 High resolution imaging

The high resolution transmission electron microscopy

(HRTEM) technique is currently one of the most informative

methods for structural analysis with atomic precision. Among

wide range of structural information, the linear defects of the

crystal lattice could be identified by this technique. The HRTEM

allows revealing of the grains and interphase boundaries

structure and measurement of the lattice distortions [98], [99],

[100].

The HRTEM image is formed by transmitted and diffracted

beams interference. Thus, the correct interpretation of the

HRTEM image contrast is not possible without usage of modern



methods of image simulation. More detail information on

formation of the TEM images could be found e.g. in [101].

1.4.7 Scanning transmission electron microscopy (STEM)

Modern transmission electron

microscopes are equipped frequently

by STEM mode attachements. In this

mode specimen is illuminated by

convergent beam and scanned similar

to the SEM technique. A schematic of

the STEM technique is presented in

Fig. 1‐16. STEM images are formed by

collecting incoherently scattered

electrons which are highly sensitive

to variations in the atomic number of

atoms in the sample (Z‐contrast

images) [102]. The scattered electrons

are recorded with a high angle

annular dark‐field detector (HAADF). So far thickness of the

samples analyzed with TEM is quite small the volume of

electron beam interacting with the sample is quite small too and

0.2 nm resolution could be achieved. The resolution could be

improved up to 0.08 nm by applying spherical aberration (Cs)

correctors. It should be noted that interpretation of STEM

images contrast is not complicated because of image contrast

proportional to the atomic number of the element. Using this

mode a significantly improvement of carbon material phase

analysis accuracy could be obtained [103], [104], [105].

1.4.8 Atomic structure simulation

Proper phase refinement and crystal orientation relative to

incident beam couldn’t be done without experimental and

simulated HRTEM images compassion. For this purpose, the

atomic models of studied nanocarbon material should be

elaborated. The atomic model may be calculated according to

known data for bulk carbon material and measured interplanar

distances and crystal orientation. On the base of the calculated

Figure 1‐16. Scheme of

scanning transmission

electron microscopy


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model and according to the TEM optical parameters the series of

simulated HRTEM images with different defocus and sample

thickness are formed. The criterion of fidelity of the phase

composition and crystal orientation determination according to

the STEM data will be a function of images correlation.

OBJECTIVES OF THIS WORK

This brief review of the carbon materials and their structural

features presents basis for motivation of the Thesis where the

following topics of scientific interest are highlighted:

to perform the theoretical analysis of the carbon atoms

contrast optimization in spherical aberration corrected

transmission electron microscope;

to analyze the atomic arrangement of guest crystal in the

inner channel of functionalized single‐walled carbon

nanotubes filled with 1‐bromoadamantane, mercury

chlorine and cupper chlorine molecules;

• to reveal the structural features of nanocrystalline and

single crystal diamond; in particular, to study of carbon

bonds transformation under high‐temperature annealing;

and to identify of structural defects in crystalline

diamond produced by CVD method;


2 General principles and

practical aspects of TEM

analysis

This Chapter will be focused of some aspects of transmission

electron microscopy techniques which are important in carbon

materials studying. In first part the theoretical principles of

electron diffraction and TEM images formation will be

discussed. The second part of the Chapter the practical methods

of electron‐optical system adjustment at spherical aberration

corrected TEM will be considered. Fine tuning of Cs‐corrected

TEM allows to achieve enhanced signal‐to‐noise ratio.

2.1 THEORETICAL ASPECTS OF ELECTRON MICROSCOPY

2.1.1 Electron diffraction

Structure analysis in transmission electron microscopy (TEM)

is based on effects occurring during interaction of electron beam

(e‐beam) with material of the sample. The periodic structure of

the sample acts as a diffraction grating, scattering the electrons.

Thus, analysis of the diffraction patterns formed by the scattered

electrons may be used for determination of the sample

structural characteristics.

In accord with L. de Broil hypothesis [106] electron has

quantum wave properties which may be described by wave

with wavelength λ, which are determined by accelerating

voltage V (i.e. by energy of electrons) by equation: ħ ħ ħ

(2.1)


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where – mass of electron, e – charge of electron, ħ – Plank

constant. For example, at accelerating voltage 100 kV the

wavelength of electron equals to 0.0037 nm, which is about

hundred times less in comparison with interatomic distances of

about 1Å (i.e. 0.1 nm) in order.

Let us consider the case of electron scattering in the sample of

material where long‐range order exists (i.e. atoms are located in

space with periodical regularity). In this case incident wave of

electron will be diffracted on atomic planes in according with

equation of Wolf‐Bragg [107]:

2 sin ɵ (2.2)

where d – interatomic distance in material of the sample, ɵ ‐ diffraction angle, n – integer. The correspondence between

structural characteristics of the material and parameters of the

diffraction pattern may be expressed by formula:

= Camera constant (2.3)

where L – distance from the sample to the plane of the

diffraction image which is registered by the camera, λ –

wavelength of electron, r – distance from the central spot

(corresponding to non‐scattered electrons) of the diffraction

pattern, and d – interatomic distance. The combination of (2.2)

and (2.3) formulas allows establishment of interatomic distance

from the diffraction pattern analysis. Further analysis may be

used for revealing of the crystal symmetry characteristics and

calculation of unite cell of the crystal. Additionally, to this

information revealing from location of the diffraction spots

analysis of their intensities allows determination of shape and

dimensions of the crystal contained in the analyzed sample.

Application of this method could be considered in example

of analysis of the electron diffraction pattern obtained from a

monolayer of close packed porphyrin−fullerene dyads

molecules (see details in Article I in the Appendix). Fig. 2‐1

shows the diffraction pattern consisting of diffraction rings that

correspond to polycrystalline structure of the sample with

randomly oriented crystalline domains. Small width of the

diffraction rings corresponds to domains size up to 10 nm.

Interpretation of experimental diffraction data is carried out, as

Chapter 2: General principles and practical aspects of TEM analysis


a rule, by applying the atomic modeling and simulation of

diffraction patterns techniques. In the Article I the analysis of

possible types of the molecules ordering in close packed

structure was made for porphyrin−fullerene dyads molecules.

Primitive cell and atoms coordinates were determined for each

possible type of the packing. Obtained data were used for

simulation of diffraction patterns using JEMS software [108].

The simulated patterns were compared with experimental data.

The best correspondence was for molecules packing into triclinic

structure with parameters of unite cell: а = 1.37 nm, b = 1.40 nm,

c = 2.47 nm, α = 89.6°, β = 90°, γ = 91°. Further details about

motivation of this work, methods for material fabrication and

other experimental conditions are disclosed on Article I (see

Appendix).

Figure 2‐1. Atomic model of the porphyrin−fullerene dyads molecule on the water sub‐phase (a). Model of the triclinic unit cell of close packed molecules (b).

Experimental (c) and simulated (d) diffraction patterns for a monolayer of close

packed molecules. Diffraction intensities distribution is shown for the simulated

pattern (d). (The images reproduced in accord with Article I – see in Appendix).


Carbon Materials


2.1.2 Principles of image formation

The combination of electron diffraction and imaging provides

the unique capability in TEM for understanding the properties

of the crystals and possible defects. TEM image is formed from

the interaction of the electrons transmitted through the

specimen. These images represent distribution of intensity of the

electron beam passed through the sample occurred due to

absorption or/and scattering in material of the sample. Similar to

diffraction patterns, the TEM images are captured using

luminescent screen, photo plate/film or CCD camera. The

example of TEM image is shown in Fig.2‐2 representing the

GaSe particles (see details in Article II from Appendix). The

TEM images demonstrate intensity distribution for e‐beam

passed through the GaSe particles located on perforated

amorphous carbon film used as a support. Intensity difference

of e‐beam may be characterized by contrast which is a ratio of

brightness of the neighboring local areas of the image. Analysis

of TEM image contrast variation allows obtaining of information

about dimensions and location of the structural peculiarities in

the sample. Relatively weak contrast for GaSe particles in Fig. 2‐

2 witnesses about small thickness of these particles which have

platelet (quasi two dimensional) shape. Other details about this

study are presented in Article II (see Appendix).

Figure 2‐2. TEM image of GaSe particles on holey carbon support film. The

images are reproduced from Article II (see Appendix).



2.1.3 TEM optical system

Formation of the diffraction patterns and TEM images in

transmission electron microscopes forms by electro‐optical

(called also as ‘optical’) system. Fig. 2‐3 schematically presents

main modes of this system. The electron gun produces a

diverging beam of electrons, which passes through aperture and

then is focused by condenser lens. Upper objective lens forms

parallel electron beam which illuminates the sample. Part of

electrons of the primary beam, which are not deflected in

sample, and electrons diffracted due to scattering on atomic

planes in material of the sample are focused then by lower

objective lens. Intermediate and projective lenses focus the

primary and diffracted beams on luminescent screen (or CCD

camera). Objective aperture allows formation of diffraction

pattern only for selected area on the sample. This scheme was

used for obtaining of the diffraction patter shown previously in

Fig. 2‐1.

The scheme of two types of TEM images formation ‐ bright‐

field and dark‐field is shown in Fig. 2‐3 (b and c,

correspondingly). In both cases sample is illuminated by parallel

e‐beam formed by the condenser lens. Bright‐field imaging (Fig.

2‐3 (b)) is obtained by focusing on detector screen (covered by

luminescent material, photo‐film/plate or CCCD camera) only

non‐scattered part of the e‐beam. Diffracted beams are rejected

in this case by objective aperture. In the opposite case of dark‐

field imaging (Fig. 2‐3 (c)) the non‐scattered part are eliminated

by objective aperture while diffracted beams are passed through

and focused on the screen.

Contrast on the bright‐field images is determined mostly by

electron‐beam absorption in the material of the sample. The

dark‐field images are formed by diffracted beams and thus

possess so called ‘diffraction contrast’. In many cases, dark‐field

imaging is very informative and use, for example, to visualize

small particles and determine their morphology.


Carbon Materials


Figure 2‐3. Scheme of transmission electron microscope with electron trajectories in

diffraction mode (a), bright‐field (b) and dark‐field (c) image mode.

To demonstrate the difference between the bright‐ and dark‐

TEM images Fig. 2‐4 shows the images of both types captured

for the area of sample containing GaSe nanoparticles on

amorphous carbon film. The dark‐field image (Fig. 2‐4 (b)) was

obtained using e‐beam corresponding to the spot, which is

denoted by circle on the diffraction pattern (see inset in Fig. 2‐4

(b)). Other details of this research are disclosed in Article II (see

Appendix).



Figure 2‐4. Examples of bright‐field (a) and dark‐field (b) TEM images of the same

GaSe nanoparticle on holey carbon support film. The dark‐field image was

obtained using diffracted e‐beam corresponding 010 spot marked on the diffraction

pattern (shown in inset to the panel) (b). See details in Article II (Appendix).

2.1.4 TEM resolution limit

The most important characteristic of TEM is, probably, its

resolution ability. The resolution itself is defined as device

ability in displaying closely located objects. A quantitative

measure of this parameter is the resolution limit of the

microscope, that is the shortest distance between two point

objects, at which they are distinguishable as separate (i.e.,

appear in the microscope as two points) (see Fig. 2‐5). The

diffraction limit ( ) is the most important value which

determines the resolution ability of the microscopes. In

accordance with Rayleigh criterion there are several

equations that have been derived to express the relationship

between wavelength and resolution [101]:

1.22 (2.4)

=.

=. (2.5)

where α is aperture angle of the lens or Numerical Aperture

(NA) of objective and condenser lens. Equation (2.5) is used

when the microscope is in perfect alignment (parallel

illumination), in this case NAcond=0. The resolution limit is

achieved at 90° opening of the objective aperture (sinα= 1).


Carbon Materials


For example, the diffraction limit would be 2.3 pm for

electron energy of 100 kV and would decrease with the

energy. Practically, the NA for electromagnetic lenses cannot

be set to 1 due to generic limitations [109]. Thus, the real

resolution of the microscopes is limited by different factors

including various imperfections of the optical system, energy

spread of electrons, instability of accelerating voltage, thermal

drift of the sample etc.

The optical systems imperfections are called lens

aberrations. The most important types of the aberrations are:

spherical and chromatic aberration, axial astigmatism, and

coma. Fig. 2‐6 presents scheme of rays in an optical system

which illustrates definitions of spherical and chromatic

aberrations.

Spherical aberration is the reason that the rays pass

through the lens are focused at different distances from the

geometrical focus of the lens. This occurs due to the

inhomogeneity of the electromagnetic field in the objective

lens. The rays passing through the lens at a distance different

from the optical axis will be undergo the various effect of the

spherical aberration which lead to deviation of rays focus

Figure 2‐5. Ray diagram representing

the lens diffraction limit. Two spots

are believed to be resolved (bottom

right) when distance between them is

more than spot radius. Focus position

may be changed from image plane

position (defocus) when lens strength

becomes weaker (underfocus) or

stronger.



position. Therefore, the focus of the lens is smeared along the

optical axis. The dimension of the focus spot determines

resolution limit of the microscope and may be estimated as

, where CS is, so called, coefficient of spherical

aberration [110]. In contrasts to the diffraction limit ( ),

which is reversaly proportional to the aperture angle of the

lens, the aberration limit is proportional to the angle aperture

in third power (Fig. 2‐5c). It follows that for any given lens of

fixed spherical aberration coefficient there should be an

optimum angular aperture which could be derived at

condition = С . Hence, we obtain:

0.61 (2.6)

Figure 2‐6. Schematic representations of spherical (a) and chromatic aberration (b)

that prevent the beam focusing at single point. The diffraction (εd) and the

spherical aberrations (εсs) limit resolution in dependence on angular aperture (α)

of the lens (c). The αopt is optimal value.

Chromatic aberration reflects the situation when the rays

of radiation with shorter wavelengths are refracted

differently than rays with longer wavelengths. It leads to

appearance of blur of focus point along the optical axis.

Chromatic aberration is a generic aberration for

electromagnetic lenses [109] and its coefficient is always

positive, i.e. electrons with higher energy are focused further

of the lens. Resolution limit due to chromatic aberration can

be expressed as [111]:


Carbon Materials


Сс Сс∆

(2.7)

where СС – coefficient of chromatic aberration, ΔE/E – relative

energy spread of electrons determined by both instabilities of

accelerating voltage and energy distribution of electrons

leaving the tip. Thus, the negative influence of chromatic

aberration can be partially overcome by using

monochromatic electron beams and highly stable accelerators.

The energy spread of electron beam is determined by type

of used electron source (or cathode). Currently two types of

cathodes are applied in TEM: thermionic electron sources,

which produce electrons as a result of thermionic emission

from tungsten filaments or crystals of lanthanum hexaboride

LaB6, heated up to high temperature; and field emission gun

(FEG), which produce electrons as a result of field emission

effect, when strong electric field is applied to the tip‐shaped

conductive material (usually made of very sharp W needle).

Typical brightness of electron beam produced by thermionic

W and LaB6 cathodes are 109 and 1010 A/m2sr,

correspondingly, while energy spread of electrons in the

beams is about 3 eV for W and 1.5 eV for LaB6 [101]. Schottky

type field emission sources provide e‐beam brightness up to

1012 A/m2sr and electron energy spread of about 0.7 eV. The

energy spread of electrons in the beam may be reduced well

below 0.1 eV by using gun monochromator. The drawback of

monochromation is a significant reduction in total beam

current, thus high brightness sources are very beneficial on

mochromated systems.

Another type of TEM optical system imperfections is

distortion of an axial symmetry of the lens which causes an

appearance of another type of aberration in focusing systems

for optical, electron and other beams. This type of aberration

is called as astigmatism. Schematic presentation of such kind

of system with astigmatism is shown in Fig. 2‐7 where it

leads to difference between the focal length for rays passing

in the plane of the figure, and for rays, which are located in

the perpendicular plane.



Figure 2‐8. Scheme of rays in the

optical system illustrating formation of

coma aberration.

Figure 2‐7. Astigmatism is the result of axial asymmetry and leads to variations

in the focal length for different angles around the optic axis. This is resulting in

co‐existence of two principle focus lines located at right angles along the axis.

Axial asymmetry of the system results in azimuthal focal

length dependence (Fig. 2‐7). In contrast to the spherical and

chromatic aberrations, astigmatism may be compensated by

special non‐round optical elements ‐ stigmators, which are

incorporated into the microscope lens to create additional

magnetic field.

Another important type

of aberrations is coma. This

aberration may appear

when incident beam is

tilted in respect of optical

system axis (see Fig. 2‐8).

Focusing of the incident

rays in different points

creates specific blurring of

the focal point with shape

of coma. Adjustment of the

optical system to avoid formation of the coma is critical for

high resolution TEM image formation.

2.1.5. Phase contrast formation and imaging with atomic

resolution

The imaging contrast which was discussed before in

Section 2.2 is formed by in‐plane variation of the amplitude

of the electron wave which is happens because of difference


Carbon Materials


in absorption and scattering of separate areas of the object.

This contrast is called ‘amplitude contrast’. Alternatively,

‘phase contrast’ in the TEM images appears as a result of

phase shift of electrons passing through the sample. The

phase shift itself cannot be detected as far the detection

system records intensity (a square of amplitude) of electron

wave. In order to convert the phase variations into the

variations of intensity an optical system should possess

particular types of geometrical aberrations, the deficiency of

focus (“defocus”) being the main used for this purpose.

The geometrical aberrations of the optical system of the

microscope may be defined via phase shift of the plane wave

after it passed throughout. Typically, these aberrations are

defined in reciprocal space as a phase shift χ dependence on

space frequency u. The space frequency u is parameter

equivalent to the reciprocal space vector which is used for

description of radiation scattering in crystals. If only defocus

and 3rd order spherical aberration are considered the

aberration function may be represented by equation [101]:

χ πΔf (2.8)

where Cs ‐ spherical aberration coefficient, Δf – defocus. The

aberration function produces main input into contrast

transfer function (CTF) of the microscope which determines

how the amplitude and phase of electron wave are

transferred from the object to the image plane.

For a weak‐phase object (thin sample) CTF is given by

[101]:

2 (2.9)

where A( ) is aperture function: A( )=1 at u≤ and A( )=0 at

u> , where is spatial frequency corresponding to the

aperture radius. The optimal shape of this function is

shown in Fig. 2‐9a. The contrast transfer function oscillates

depending on space frequency. The phase contrast can be

maximized by maximizing the integral below CTF. As far as

Cs is fixed (for non‐corrected microscopes, see consideration

for corrected case below) this optimization is done by



varying the value of Δf. The optimal value for Δf is given by

the formula [109]:

(2.10)

and is commonly called Scherzer focus. Due to the fact that

Cs is generically positive for electromagnetic lenses, Δfopt is

negative and as a consequence CFTfopt has broad band path at

negative values. Practically this means that optimal phase

contrast will be negative, i.e. atoms in TEM image will

appear dark on light background.

The equation (2.8) describes an ideal case with constant

aperture function A(u) while in a real microscope the transfer

function is attenuated according to equation [101]:

(2.11)

where с and are, respectively, the damping envelopes

due to chromatic aberration (Cc) and beam convergence (α):

exp (2.12)

exp (2.13)

where δ is defocus spread caused by fluctuation of

accelerating voltage (ΔEv), the energy spread of the electrons

(ΔEe), emitted by the filament, and the energy loss of

electrons (ΔEs) due to inelastic scattering in a specimen

generated by variations of wavelength, resulting in chromatic

aberration (Cc). Additionally, to that variations of current (ΔI)

of the objective electron lens lead to the focal length change.

Defocus spread (δ) can be expressed as [101]:

δ= Сс 2 (2.14)

The resulting CTF function of the microscope Ttotal(u) may

be presented graphically as it shown in Fig. 2‐9b. To exclude

effect of phase variation the images should be obtained in the

range of space frequencies (u) spans from zero coordinate

point to first intersection point of the function Ttotal(u) and

horizontal axis. This point defines point‐to‐point resolution

(pp) and corresponds to a minimum distance between two

point‐like features in the object, which can be distinguished

from each other:


Carbon Materials


0.65 (2.15)

Figure 2‐9. The ideal CTF function of microscope without taking into account

chromatic aberration and beam convergence dumping envelops (a). Dependence of

with respect to u modified by the damping envelope (blue dashed

curve) – (b). The calculations are made for Δf=‐100nm, Cs=2.2 mm.

Spatial frequency at which Ttotal(u) becomes zero defines so

called information limit (marked at Fig. 2‐9 (b) as ). This

point determines the highest frequency which can be

transferred through the optical system:

(2.16)

The information limit is far beyond limits of point

resolution for the microscopes with FEG sources because of

their high space and time coherency. The high resolution

atomic scale imaging may be achieved but the complex

dependence of the contrast on space frequency requires

usage of image simulations for correct determination

positions of atoms.

Fig. 2‐10 shows an example of the high resolution imaging

with TEM. This image presents atomic structure of thin



quasi‐two‐dimensional Si crystal. The distances and angles

between atomic columns were determined using Fourier

transformation and diffraction patterns simulations.

Estimation of the distances between spots in the diffraction

patterns and angles between vectors directed from the center

to the analyzed diffraction spots allows establishing of cubic

primitive cell and determining the orientation of the crystal in

respect to the incident beam of electrons. To confirm

correspondence of the experimentally obtained image to

silicon, the methods of computer simulations were used to

obtain images with atomic structure and the simulated

images were compared with the experimental data. Inset in

Fig. 2‐10а’ shows calculated images for silicon crystal with

thickness of 6.9 nm at defocus 65.5 nm. Inset in Fig. 2‐10b’

shows the calculated image for the crystal with thickness of

9.3 nm at defocus 64.5 nm. More details about this study may

be found in Article III (see Appendix).

Figure 2‐10. HRTEM image of the flat Si crystals. The diffractograms of the

selected areas (a) and (b) correspond to [101] (c) and [112] (d) zone axis,

respectively. Image simulation (insets) are shown on filtered image correspond to

crystal thickness of 6.9 nm and defocus of 65.5 nm (a’); and for crystal thickness

9.3 nm and defocus is 64.5 nm (b’). See further details in Article III (Appendix).

2.1.6 Improvement of resolution by aberration correction

Resolution of the microscope can be significantly

improved by correction of one of the most important


Carbon Materials


parameter – the spherical aberration of the objective lens. For

the first time the method of the aberration corrections has been

proposed by Sherzer in 1947 [112], who has developed strategies

to correct the chromatic and the spherical aberration of a round

magnetic lens. In one of proposed variants the corrections

include introduction of non‐round optical elements into

magnetic lens system. In this case the magnetic field of reduced

azimuthal symmetry is produced due to rotational asymmetry

of the lenses. Efficient correction of the lens aberrations may be

achieved, potentially, with such multi‐pole lenses. In 1990 the

first aberration corrector with a new hexapole correction system

has been proposed by Rose [113]. These correctors have proved

possibility for increasing of resolution of electron microscopes.

Figure 2‐11. Schematic presentation of the spherical aberration corrector with

hexapoles.

The compensation of the spherical aberration for STEM mode

has also been started in the 1990s. The proof of principle has

been demonstrated in 1997 by Krivanek [114]. At now days two

designs of spherical aberration corrector are in use: the

octupole/quadrupole type and the hexapole type. A schematic

representation of the hexapole corrector is shown in Fig. 2‐11.

Corrected system is thus composed of the objective lens, the

telescopic transfer doublet, and the hexapole corrector itself. The

corrector consists of two sextupoles (hexapoles) and another

telescopic round lens doublet formed by two identical round

lenses separated from each other by a distance of double focal

length. The doublet reverses the course of the paraxial

trajectories images the sextupole located at the front‐focal plane



of the first lens with magnification onto the second sextupole

placed at the back focal plane of the second lens.

For aberrations correction, it is necessary in the first stage

to measure aberrations and then to compensate them. There

are different approaches for the aberration estimation,

including: applying correlation functions for simulated and

experimental images; estimation of the reconstructed wave

functions; and diffractograms analysis. The most widely used

approach is the diffractograms analysis of the amorphous

carbon film including Fourier transformation of high

resolution TEM image. This method was proposed at the first

time by Thon [115]. An example of diffractogram of

amorphous film with 23 nm defocus is represented in Fig. 2‐

12. It consists of series of light and dark rings and their

intensity may be expressed by equation (compare to 2.8):

| | ~sin (2.17)

From this dependence one can see that dark and light rings

of the diffractogram correspond to bands and gaps of CTF.

The values of defocus, two‐fold astigmatism, spherical

aberrations, and even high‐order aberrations may be

calculated on the basis of analysis of the diffractograms using

contrast transfer function which is determined by the optical

system aberrations.

Small two‐fold astigmatism may be calculated using

measured values for diameters of light and dark rings which

will have elliptical form (see Fig. 2‐12d). The light rings

correspond to sin(χ(u))= 1, and the dark rings correspond to

sin(χ(u))=0. Thus, we obtain two values:

for dark rings χ u , where n is odd, (2.18)

for light rings χ u , with even n. (2.19)

Substituting these values into equation for aberration

function (2.7) we obtain:

2 (2.20)

For defocus estimation, it is better to use first dark rings than

first light ring. For the first dark ring n=2. The second term in

(2.20) can be neglected for the microscopes with spherical


Carbon Materials


aberration correction because of small value of the spherical

aberration coefficient (tens of μm). Thus, we obtain:

(2.21)

By estimation of diameter of first ring in two perpendicular

directions (u1 and u2 in Fig. 2‐7d) we can calculate defocus and

astigmatism values:

, (2.22)

The resulting defocus and two‐fold astigmatism may be

calculated as average and difference of these two parameters:

(2.23) | |

(2.24)

Figure 2‐12. (а) – Profile of relative intensity distribution created from the center of

the diffraction pattern along arrow shown in panel c. (b) ‐ CTF function of the

microscope (300 kV, Cs=1.2 mm, С1=23 nm). (c) and (d) – calculated diffraction

patterns for double‐axis astigmatism A=0 nm and A=81 nm, correspondingly.

The estimation of spherical aberration coefficient is more

complicated. For uncorrected microscopes one should take it

into account dependence of Cs versus positions of rings in

equation (2.19). By plotting nu‐2 versus u2 one obtains the

straight line with slope Csλ3 [116].

For Cs corrected microscopes usually aberration estimations

carried out by analysis of diffractograms obtained at tilted

illumination. This method was proposed in 1978 by Zemlin and



is realized using, so called, Zemlin tableau. It provides more

accurate aberrations measurement [117]. Moreover, it allows

calculation of higher‐order and non‐centrosymmetric

aberrations. The Zemlin tableau example is shown in Fig. 2‐13.

The diffractograms were obtained under beam tilting (with

azimuthal rotation from 0 to 2π) and along optic axis (central

image). Fig. 2‐13(b), (c) shows diffractograms before and after

aberration compensation. Aberration compensation leads to

ellipse‐to‐rings shape transformation (Fig. 2‐13c).

Figure 2‐13. Zemlin tableau: series of diffractograms obtained at tilted electron beam

(a). Examples of diffractograms before (b) and after (c) aberration correction.

Cs‐correction allows reduction or even conreverce the

value of spherical aberration. Comparison of the CTF

functions at Sherzer defocus for uncorrected (Cs=1.2 mm)

and corrected (Cs=0.02 mm) TEM with accelerating voltage of

300 kV are presented in Fig. 2‐14. The point resolution is

drastically improved from 0.22 nm to 0.08 nm. The

simulations are made using CTF Explorer software [118].

Thus, Cs‐correction allows control over the value of Cs in

addition to ∆f, thus extending the parametric space available

for optimizing the imaging. At these conditions the

aberration function should be reviewed. The (2.7) equation

was derived with assumption that Cs is `dominantʹ and fixed

geometrical aberration. The role of higher order aberration


Carbon Materials


was assumed to be negligible. To describe the information

transfer for ultra‐high resolution imaging besides the correctable

twofold axial astigmatism, additional anisotropic axial

aberrations, such as threefold astigmatism and axial coma,

would need to be considered[119,120].

Figure 2‐14. Comparison of the contrast transfer functions at Sherzer defocus for

uncorrected and corrected TEM with accelerating voltage of 300 kV: (a) ‐ Cs=1.2

mm, (b) ‐ Cs=0.02 mm.

The aberration function will have a more complex view

taking into account higher order and asymmetric geometrical

aberrations. Using a common notation [121,122], it could be

represented as:

12

12

13

14

14

… (2.25)

where complex scattering angle , and its complex

conjugate is . The notations of geometrical axial aberrations



used from now on are represented in Table. 2.1. The detail

review of theory of aberration is presented in [122].

Table 2.1. Geometrical axial aberrations

Aberration Symbol Defocus (∆f) C1 Two-fold astigmatism (A) A1 Second-order axial coma B2 Three-fold astigmatism A2 Third-order spherical aberration (Cs) C3 Third-order star aberration S3 Four-fold astigmatism A3 Fourth-order axial coma B4 Fourth-order three-lobe aberration D4 Five-fold astigmatism A4 Fifth-order spherical aberration C5 Six-fold astigmatism A5

So, for Cs‐corrected microscopes the estimation of

aberration for it proper adjustment is necessary.

2.2 OPTIMIZATION OF SPHERICAL ABERRATION CORRECTION FOR NANO-CARBON MATERIALS (EXPERIMENTAL PART)

Significant improvement of spatial resolution and ability for

study of the samples at low accelerating voltages become

possible with application of the spherical aberration correctors.

Low accelerating voltage allows increase image contrast, to

reduce the radiation damage (which is important for radiation

sensitive materials), increase of sample stability and other

important enhancements. All these advantages of structural

characterization at low accelerating voltage are especially

important for low atomic number materials such as nano‐

carbons [123].

As it was mentioned before, the possibility to vary C3 in

addition to ∆f gives the opportunity for further optimization of


Carbon Materials


CTF in order to compensate untunable fifth‐order spherical

aberration C5. The C5 parameter is the next after С3 in its

contribution into aberration function (2.25). Currently C5

correctors are still in the development stage and are not in

common use, so we will assume that the value of this aberration

is fixed by the construction of the optical system. The criteria for

phase contrast CTF optimization is essentially the same as was

mentioned before for uncorrected systems: the integral under

sin should be maximized. This is achieved by a proper

choice of the combination of defocus C1 and third‐order

spherical aberration C3 in order to compensate the action of

fifth‐order spherical aberration C5.

There is substantial number of works [124–126] discussing

optimal settings for Cs‐corrector. Typical factors, which are

accounted for are the values of C5 [121,127,128], chromatic

aberration [113,119,125], contribution of non‐linear terms in CTF

[117,129], contribution of amplitude contrast [122]. It is a

common perception that negative spherical aberration imaging

(NCSI) [130,131] in preferential for the corrected systems.

However, this was so far demonstrated for strongly scattering

crystals, and up to our knowledge was not confirmed

theoretically for light materials like carbon. There are a number

of other issues related to carbon materials, which should be

considered as additional parameters contributing to

optimization, namely: necessity to maximize the contrast in

order to reduce the dose, necessity to use sub‐threshold

accelerating voltage (≤80kV), stringent constrains for accuracy of

aberration correction at low accelerating voltages to mention a

few. These considerations lead us to perform an extensive

theoretical study on optimization of the contrast of carbon

atoms in corrected system and in particular the performance of

the corrector itself.

2.2.1 Equipment

The measurements in this work were performed on a high‐

resolution transmission electron microscope Titan 60‐300

TEM/STEM (FEI, The Netherlands) with high‐brightness field



emission gun X‐FEG at an accelerating voltage of 80 kV (Fig.2‐

15). To minimize the contribution of the chromatic aberration,

the microscope was equipped with a monochromator, which

allows reducing energy beam spread less than 0.1 eV. The

spherical aberration of the objective lens was compensated by

CEOS (CEOS, Germany) corrector. HRTEM images were

acquired by Gatan (Gatan, USA) UltraScan1000 2k x 2k CCD

camera.

2.2.2 Test samples

The graphene was chosen as a test sample because it has one

atom thickness, which eliminates unambiguity in thickness

determination and allows for direct interpretation of the images

in terms of atomic structure. Graphene samples were grown by

chemical vapor deposition method using 25 μm copper foil as

the substrate. The graphene monolayer is transferred onto TEM

Au Quantifoil grids with holey carbon film using

polymethylmethacrylate (PMMA) as the sacrificial polymer

layer and ferric chloride as the copper etching agent [132].

Figure 2‐15. Transmission electron microscope FEI Titan 60‐300 with X‐FEG gun,

monochromator and CEOS Cs‐spherical aberration corrector.


Carbon Materials


2.2.3 Procedure of aberrations measurement and correction

The screenshot of CEOS Cs‐corrector interface is presented in

Fig. 2‐16. Correction procedure consists of two main parts: first

aberration evaluation and second correction of them. Aberration

estimation is carried out by acquiring Zemlin tableau [117]. The

process takes place in an interactive mode, where the role of the

operator is to select the optimum measurement conditions

(magnification, beam tilt angle, exposure time, defocus),

measuring and evaluating the correctness of the decision to the

subsequent step. The aberration measurements and correction

are made by corrector software.

2.2.4 Roadmap for corrector optimization

Following steps were performed to achieve the purpose of

this Chapter:

Evaluation the aberration measurements algorithm,

revealing systematic errors and accuracy;

Determination of the optimal measurement

conditions (beam tilt angle, magnification) for better

accuracy;

Determination the optimal combination of C1, C3 and

C5 parameters for maximum contrast of carbon atoms

(maximal signal‐to‐noise ratio);

Experimental evaluation of the findings on graphene

sample.



Figure 2‐16. Layout of Zemlin tableau measurement.

2.2.5 Evaluation of corrector software algorithm

For evaluation of CEOS software aberration measurements

algorithm the series of calculated images with specified

aberrations are simulated. These calculated images are then

analyzed by the corrector program, which measures these

aberrations value. Comparison of simulated and measured

values gives the estimation of corrector performance.

Emulating of Zemlin tableau requires simulation of the

images at tilted illumination with account for off‐axial

incoherent aberrations, in particular for off‐axial chromatic

aberration. In contrast to on‐axis case, where chromatic

aberration can be accounted for as the blur of defocus and thus

as in‐plane symmetric smearing of the image, at tilted

illumination this blur of defocus is directional (see fig.2‐17 left).

In practice this aberration is seen on the FFTs of high resolution

images as so called “bananas”. An example of (simulated) image

of amorphous film at tilted conditions and corresponding FFT is

represented at fig.2‐17 middle and right respectively. Simulation

of this phenomena so far was not demonstrated, though there

are a few papers discussing CTF for tilted illumination for a

very simplified cases (not accounting for the sample, dose,

camera, aberrations higher than defocus, etc.) [125,133]. Here we

propose and realize the method of simulating tilted illumination


Carbon Materials


with account for the sample properties (thickness, density,

composition, structure, inclination), all axial aberrations, as well

as imaging parameters (magnification, dose, CCD MTF). The

method uses a direct integration of images through

(2.26)

assuming Gussian distribution of energy spread. An example of

thus calculated Zemlin tableau is represented at fig 2‐18.

Figure 2‐17. (a) – Scheme of considered tilted illumination geometry for simulating

the HRTEM image of amorphous carbon film (b) and it FFT (c).

The data set of 17 calculated diffractorgams calculated from

simulated images with illumination inclination of 18 mrad is

shown in Fig. 2‐18.). Calculations are done for every tilt angle (τ)

from 10 to 35 mrad with 1 mrad step (Fig2‐18, left). For each tilt

angle, the set of diffractogram is calculated. Thus, around 600

diffractograms are calculated and analyzed. All sets of

simulated data are loaded into corrector software and aberration

value are estimated.

Simulated data are compared with experimental

measurements on graphene sample. As well as to the calculated

data, aberration measurements are carried out in the

illumination tilt range from 10 to 35 mrad.

Measured C3 value versus electron beam inclination angle for

the calculated and experimental data are shown in Fig. 2‐19. The

experimental points are shown on the plot by blue squares, the



theoretical values shown by crosses. Analysis of obtained data

by fitting curve allows evaluation C5 parameter.

Figure 2‐18. An example of tableau calculated from simulated images which include

aberrations up to C5 (beam tilt angle τ=18 mrad). Nominal defocus C1 is ‐250 nm,

C3= ‐30 μm, C5=12 mm, 500 nm of coma is introduced in addition.

We have calculated such tableau for the tilt ranges (τ) from 10

to 35 mrad assuming the following settings: amorphous carbon

sample thickness 1nm, sample temperature 300K, accelerating

voltage 80kV, C3=‐30um, C5=12mm, defocus 250 nm, defocus

spread 10nm, magnification x250K. All sets of simulated data

are loaded into corrector software and Cs‐values and its

“confidence” levels are measured. The results of these

measurements are represented on Fig. 2‐19.

Defocus value measured by corrector is right (250.1 nm),

however the values of C3 significantly deviate from the

simulated value (horizontal dashed line, Fig. 2‐19). Error bars of

the measured points represent confidence levels reported by the

corrector software and are order of magnitude smaller than the

deviations of C3 from the true value. Observed dependence of C3

from the tilt range could be potentially explained by neglecting


Carbon Materials


the contribution of C5 in fitting algorithm (which is most

probably the case). In this case the measured value of C3 would

deviate from true value (for the type of tableau used) as:

С 1.576 (2.27)

This curve calculated for simulation parameters is shown in

green on Fig2‐19. It is seen that even after accounting for C5

there are substantial deviation of the measurements from the

real data. The closest (and precise) measurement is achieved at

18 mrad with a systematic shift from the true value of about

+7 um. In the further practical work, we will be using this value

for measurements as well as will be accounting for this

systematic shift in C3.

Independent calculations (not shown here) revealed

optimum magnification, optimal defocus, signal‐to‐noise ratio

necessary for reliable measurement as well as potential errors

induced by sample tilt, sample thickness and calibrations

imperfection.

2.2.6 Contrast transfer function adjustment

As has been discussed above CTF optimization consists of

estimation of optimal combination of defocus C1 and third‐order

spherical aberration C3 providing the maximum integral under

the CTF. This achieved by shifting the high frequency cutoff of

CTF to maximum possible value. Two limiting factors can be

considered in this respect: cut‐off due to incoherent envelope

and cut‐off due to phase contrast oscillations caused by

uncorrectable C5 contribution. These two cases can be

discriminated by comparing contributions of one and another

shown in Fig. 2‐20. Presented here are the dumping envelope

and phase shift induced by C5 calculated for the parameters of

our microscope (80kV, Cc=1.7 mm, C5=12.4 mm, ∆E=90meV). It

is seen that dumping envelope still allows a substantial contrast

transfer at 0.1 nm, while the phase shift induced by C5 crosses

π/2 value at about 0.2 nm, meaning from this frequency CTF will

start oscillating. Thus, we can conclude, that our imaging is C5

limited and we will aim to optimize C1 and C3 for compensating



C5 (see also [127]). Such optimization results in the values

C1=10 nm, C3=‐24 μm assuming C5=12.4 mm.

Figure 2‐19. Dependencies of coefficient С3 on deviation angle of the beam in respect

to the optical axis. The crosses – data obtained by analysis of the calculated

diffractograms. Dashed line shows the value implemented in simulations. Green curve

represents the correction curve accounting for C5 (see text and eq. 2.25).

Figure 2‐20. Two limiting factors of Cs‐corrected microscope: incoherent envelope

and phase contrast oscillations caused by uncorrectable C5 (calculated for the

microscope parameters: 80kV, Cc=1.7 mm, C5=12.4 mm, ∆E=90 meV).


Carbon Materials


We will demonstrate the importance of perfect alignment of

correcting system on the base of the following examples:

In the first case CTF is treated in a “classical” way, i.e. the

contribution of C5 is neglected and optimum C1 and C3 are

found in a classical Scherzer sense. Such approach is still

common in scientific literature (see for example [134]). At the

same time the value of C3 reported by corrector is trusted in

addition. The consequence of such approach is illustrated on Fig.

2‐21 (a‐c), where (a) is the CTF the operator expects to have, (b)

is the true CTF and (c) shows how this reflects on the image of

graphene dislocation core.

In the second case the operator is smart enough to account

for C5, though is not sufficiently smart to be concerned about

corrector measurements. In this case the value of C3 set up for

imaging will differ but systematic offset discussed above, which

will result in deviation of CTF from expected shape (Fig. 2‐21 d‐

e) and corresponding disturbance of the image (Fig. 2‐20f).

Only complete accounting for system aberrations and

corrector systematic errors allows forming a perfect CTF shape

(Fig. 2‐20 g‐h) ensuring clear structural image (Fig. 2‐20 i) and

maximum contrast which represented on relative intensity

profile (Fig2‐20 j). Proper setup of conditions thus allows to

increase the contrast (and thus to decrease necessary dose) by

almost 30%.

An example of experimental images of a single wall carbon

nanotube and graphene were obtained using optimized

corrector adjustments are shown in Fig. 2‐22.



Figure 2‐21. Comparison of different CTF optimization and simulation of HRTEM

images of graphene in accordance with optical system parameters. (a‐c) In the first case

CTF is treated in assumption of neglected C5 and optimum C1 and C3 are found. (a) is

the CTF the operator expects to have, (b) is the true CTF and (c) simulated HRTEM

image. In the second case C5 is considering in accordance with measured by corrector

software value. (d) and (e) expected and real CTF and (f) simulated HRTEM. Last case

the complete accounting for system aberrations and corrector systematic errors.

Coincide of expected (g) and real (h) CTF shapes and corresponding HRTEM image (i).

Profile of relative intensity distribution (j) obtained along lines selected in HRTEM

images (c, f, i).


Carbon Materials


Figure 2‐22. Examples of experimental HRTEM images for single wall carbon

nanotube (a) and for single layer graphene (b) obtained after optimization of the

microscope adjustments.

SUMMARY OF THE CHAPTER 2

Analysis of aberration measurement algorithm of CEOS

software revealed the presence of a systematic error due to the

absence of fifth‐order spherical aberration coefficient in fitting

algorithm. Dependence of the aberration evaluation accuracy

with respect to standard beam tilt angle was determined. The

beam tilt angle about 18 mrad corresponds to the most correct

optical system aberration evaluation. Analysis of contrast

transfer function of microscope with С5 optimization in different

assumptions reveals optimized defocus and C3 parameters for

highest signal‐to‐noise ratio.


3 Atomic arrangement of

guest crystal in the inner

channel of functionalized

single‐walled carbon

nanotubes

The functionalization of the carbon nanotubes leads to

significant change of nanotube electronic properties [26,76].

Doping of SWCNT films with inorganic materials can increase

the film conductivity and change the optical absorption [76,135–

137]. Besides SWCNT could be used to create novel one‐

dimensional structures inside the nanotube channels, which are

unstable as freestanding systems [138–140]. Among the various

functionalization nanotubes methods, the filling by guest crystal

molecules is the most effective one [43,73]. Properties of the

nanocomposite based on filled nanotubes can be correctly

interpreted only if structural organization of internal channel of

the nanotube is revealed [47,77]. Among different structural

methods high‐resolution transmission electron microscopy is a

leader due ability provide local analysis and high spatial

resolution [78,141]. These parameters are very essential for

studying crystals with size of few nanometers and rarely

possessing of high order of atomic arrangement. Application of

chemical analysis techniques in TEM allows revealing

composition and the real crystal structure of the encapsulated

crystal [47,78,101].


Carbon Materials


In this Chapter, we will consider different structural

organization of the guest crystal molecules in inner channel of

filled SWNTs:

SWCNT filled by 1‐bromoadamantane molecules

(1‐bromoadamantane@SWCNT)

SWCNT filled by mercury chloride molecules

(Hg2Cl2@SWCNT)

SWCNT filled by copper chloride molecules

(СuCl@SWCNT)

3.1 SINGE WALLED CARBON NANOTUBE FUNCTIONALIZED BY ADAMANTANE MOLECULES

CNTs may be used as a nanoreactor for producing of new

one‐dimensional crystals. One of the most interesting tasks is to

grow one‐dimensional diamond crystal. One of the approaches

to reach this goal is to analyze the 1‐bromoadamantane

molecules packing into single walled carbon nanotube channel.

It is supposed that after CNTs filling procedure and heat

treatment these molecules will crystallize into diamond

structure. Control of filling results by spectroscopic techniques

is difficult due impossibility to detect the presence of carbon‐

containing molecules inside carbon nanotubes. Due to presence

of bromine atom in the guest molecule filling results could be

analyzed by EDS technique.

The SEM image of the filled SWCNTs film and STEM image

of the nanotubes bundle are presented in Fig. 3‐1. EDX spectrum

obtained from the selected area (white rectangle) is clearly

revealed the presence of bromine Fig. 3‐1c. Bromine K‐ and L‐

lines were detected for 1‐bromoadamantane@SWCNT. Low

signal from Fe, which can be attributed to the presence of

catalytic particles in the samples, is also observed.

Structure of filled nanotubes was studied by high resolution

electron microscopy technique. The carbon structures, which

looks like deformed nanotube in the inner SWCNT channel,

were revealed. A series of HRTEM images with 0.1 sec exposure

Chapter 3: Atomic arrangement of guest crystal in the inner channel of

functionalized single‐walled carbon nanotubes


were recorded (Fig. 3‐2 (a)‐(e)). This technique was applied for

visualization of free bromines atoms, which were detected as

bright spots (marked by arrows).

Figure 3‐1. SEM (a) and STEM (b) images of 1‐bromoadamantane@SWCNT. An

EDS spectra from nanotubes bundle (selected by white box in (b)) – (c).

Figure 3‐2. Series of HRTEM images of 1‐Bromoadamantane@SWCNT for

visualization of free bromines atoms and carbon structures moving inside the carbon

nanotube (a‐e). Experimental (white box in (b)), simulated HRTEM images and

atomic model of SWCNT with Br atom (f), (g), (l). The relative intensity profiles

through the marked line for experimental and simulated images (h) and (k). (The

images reproduced in accord with Article IV – see in Appendix)


Carbon Materials


Analysis of series of HRTEM images revealed that bromine

atoms migrate inside the tube during recording. Some of them

vizualized as free‐standing atoms or embedded in the carbon

structures. For proper contrast interpretation of the HRTEM

images, the simulations were performed for filled SWCNT and

for free bromine atom into accordance the TEM optical

parameters (Fig. 3‐2g). The relative intensity profiles through

the marked line (red line) from simulated and experimental

images are shown in Fig. 3‐2 (h)‐(k). Free bromine atom inside

SWCNT can be easily identified by high contrast difference due

to their low atomic number. Good agreement is observed for

experimental and calculated profiles.

Thus, it was shown that functionalization of SWCNT by 1‐

bromoadamantane leads to formation in inner nanotube carbon

structure and free standing bromine atoms. Other details are

presented in Article IV (see Appendix).

3.2 STRUCTURE CHARACTERIZATION OF HG2CL2 CRYSTALS LOCATED IN THE INNER CHANNEL OF SWCNT

Another example of functionalized carbon nanotubes studied

in this work is SWCNT filled by mercury chlorine molecules.

The molecules were loaded inside channels of single‐walled

CNTs using melting‐phase filling procedure, which provide

larger encapsulation yield. Powders of SWCNTs and HgCl2

were heated at temperature 290° C which higher than HgCl2

melting temperature by 17° C. The heat treatment was carried

out during 16 hours.

On the first step the control of the filling results was

reformed by scanning transmission electron microscopy and

EDS techniques. Due to large difference of carbon, mercury and

chlorine atomic numbers HAADF STEM image with Z‐contrast

clearly discriminate filled (bright contrast inside nanotube) and

empty tubes. Typical STEM images of the filled SWCNTs are

presented in Fig. 3‐3. About 80% of tubes were filled that




indicates good outcome of filling process. Encapsulated crystals

with length varying from 1 to 150 nm were uniformly

distributed in the channels of SWCNTs. Chemical composition

was determined by EDS analysis. The obtained spectra shown

peaks of mercury, chlorine, carbon, copper and oxygen (Fig. 3‐

3c). Copper signal is a common artifact because used TEM

support grid is made of copper; oxygen is, most probably,

related to oxygenation of carbon nanotubes during sample

exposure in air.

Figure 3‐3. STEM images with various magnification of Hg2Cl2 filled SWCNT (a), (b).

EDS spectra from filled SWCNT rope (c).

The structure of mercury chloride nanocrystals was

evaluated on the base of a set of series of HRTEM images. Two

distinct crystal orientations were observed and analyzed.

Diameter of SWCNTs evaluated by unfilled region of

corresponding nanotubes is about Dm= 1.7 nm. The {110}

graphene reflections were used as internal calibration standard.

Applying the Fast Fourier Transform (FFT) the 2D lattice

parameters of encapsulated crystal and angles between

reciprocal vectors were measured. In Fig. 3‐4 and Fig. 3‐5 are

presented HRTEM images with various encapsulated crystal

orientations. A periodicity of encapsulated crystal d1 was

determined to equal to 0.271 nm and d2 = 0.248 nm (the ratio

d1/d2 = 1.09), an angle α = 88°. A periodicity is S1 = 0.269 nm, S2 =

0.324 nm, the ratio S1/S2 = 0.83, an angle α = 85°.


Carbon Materials


Experimental data were compared with crystallographic

parameters for HgCl2 (ICSD 23277) and Hg2Cl2 (ICSD 65441)

phases of bulk crystal. Comparison of calculated diffraction

patterns with experimental FFTs shows the best match for

Hg2Cl2 phase.

Figure 3‐4. HRTEM image of Hg2Cl2 filled SWCNT (a). Interplanar distances and

angles between reciprocal vectors were evaluated by FFT form selected region of

nanotube (b). Simulated diffraction patterns for Hg2Cl2 crystal with [100] zone axis (c).

Comparison of filtered HRTEM image (d) of encapsulated crystal with simulated one

(e) which calculated in according to proposed atomic model (f).

Crystal orientation was evaluated as [100] Hg2Cl2 zone axis

(Fig. 3‐4c) and as [110] Hg2Cl2 zone axis (Fig. 3‐5c).

The atomic model of filled nanotube was proposed in

accordance with the identified structure and orientations of

encapsulated crystal. A fragment of (13, 13) SWCNT with

diameter 1.7 nm corresponded to measured one from the

HRTEM images was taken for simulation. Encapsulated crystal

is oriented [001] Hg2Cl2 direction along to the long tube axis.

Calculated HRTEM images simulated according to atomic

models of filled SWCNT was in good agreement with

experimental images. Thus, it could be concluded that




encapsulated crystal structure identified on base of HRTEM

data is correct.

To summarize, we observed that stochiometry of mercury

chloride crystals changes from HgCl2 to Hg2Cl2 during the guest

molecule loading into the inner channel of S2WCNTs.

Encapsulated crystals were oriented [001] Hg2Cl2 direction

parallel to nanotube long axis. This study provides new

information about chemistry inside the carbon nanotubes and

synthesis of new functionalized nanocarbon material.

Figure 3‐5. HRTEM image of Hg2Cl2l@SWCNT (a). FFT form selected region of

nanotube (b). Simulated diffraction patterns for Hg2Cl2 crystal with [110] zone axis (c).

Comparison of filtered HRTEM image (d) of encapsulated crystal with simulated one

(e) which calculated in according to proposed atomic model (f).

3.3 CRYSTAL STRUCTURE OF 1D CUCL@SWCNTS

Functionalization of nanotubes by gas phase filling is one of

widely applied technique. During the filling procedure

treatment of nanotube almost does not damage the structure of

the media. Gas‐phase approach leads to much cleaner SWCNT

surface comparing with a liquid method of filling. This filling

method was used for loading the CuCl molecules into SWCNT


Carbon Materials


channel. It was shown that this functionalization leads to a

significant increase in transmittance of the initial nanotube films.

We shown that this type of filling allows improvement of optical

properties of the macroscopic samples of SWCNTs (for instance,

films) without damaging their crystalline structure.

Transmission electron microscopy was applied for structure

characterization of the film of filled SWCNTs. TEM analysis of

the SWCNT sample filled by copper chloride shown that

nanotubes were predominantly bundled together into ropes.

Estimation of the yield of the nanotubes filling was performed

by STEM and EDS methods. The CuCl nanocrystals were clearly

distinguished by HAADF STEM imaging with Z‐contrast (Fig.

3‐6a). The EDS spectra acquired from the filled tubes shown

peaks corresponding to gold, copper, chlorine, oxygen, and

carbon (Fig. 3‐6b). Gold signal in this case is explained by usage of golden TEM support grid. Oxygen was most probably related

to oxygenation of carbon nanotubes during procedures used for

opening of nanotube ends. Besides a very low signal to noise

ratio the presence of both of Cu and Cl in the tubes can be

concluded without doubt.

Figure 3‐6. STEM image of CuCl@SWCNT (a) and EDS spectra (b) from selected on

(a) area.

Formation quasi one‐dimensional CuCl crystals inside the

SWCNT inner channel was revealed by HRTEM. The {110}

graphene reflections from SWNTs was used for precise

magnification calibration as internal standard. The diameter of




SWCNTs was evaluated on unfilled region of nanotubes as

Dm=1.7 nm. Several crystal orientations were observed and

analyzed. The projection of CuCl crystals observed most often in

the CNT sample is shown at Fig. 3‐7a. The interplanar distances

and angles between the reciprocal vectors were analyzed using

fast FFT method (Fig. 3‐7b). The obtained values are: d1=0.33 nm,

d2=0.33 nm, α≈60o. The interplanar distances and angles for

second orientation were determined as follows: d1=d2=0.19 nm,

α≈90o (Fig. 3‐7d).

Experimental data were compared with crystallographic

parameters for known polymorphs of copper chloride. The best

agreement was achieved for CuCl described by zinc blende type

structure with cell parameter a = 0.541 nm [142]. Comparison of

CuCl interplanar distances with experimental data shown a

good agreement for [110] CuCl projection. The main motif in

this projection is similar to observed “open hexagon”. But to

precise fitting to observed data some lattice distortions should

be supposed. CuCl lattice could be represented by two subcells

superposition. Due to integration of nanotube wall with CuCl

subcells, the ~8% subcell displacement could be assumed. As a

result, the unit cell symmetry is decreased: cubic type is

transformed to rhombohedral type. According to this

assumption the atomic model of CuCl filled nanotube was

proposed (Fig. 3‐8). Simulation of HRTEM images was

performed for this model. As could be seen from Fig. Fig. 3‐8 the

rotation of encapsulated crystal around long tube axis at 90°

allows the cubic motif which also was observed experimentally.

The coincidence of the experimental and calculated HRTEM

images for two crystal orientation confirm correctness of the

proposed atomic model.

To summarize, the functionalization of single walled carbon

nanotube by filling with CuCl molecules leads to quasi 1D

crystal formation. Applying of HRTEM techniques allows to

identify the crystal structure of encapsulated crystal. Loading of

copper chloride molecules into the nanotubes during gas‐phase

synthesis leads to crystal lattice transformation from cubic to

rhombohedral symmetry types.


Carbon Materials


Other details about CuCl functionalized single walled

carbon nanotubes are presented in Article V (see Appendix).

Figure 3‐7. HRTEM images of encapsulated 1D CuCl@SWCNT with two typical

orientations are presented in (a), (c) and corresponded Fast Furrier Transformation (b),

(d). The main motif of atomic arrangement is represent by the models in the middle.

Figure 3‐8. Atomic model of CuCl filled SWCNT viewed in two orientations and

simulated HRTEM images.





The results presented in Chapter shown that high‐resolution

transmission electron microscopy is a powerful method for

carbon materials study on atomic level. It allows control of

nanotube filling process, structural peculiarities of guest

molecules inside the tubes and provides an information about

structural transformation during synthesis.

In particular, we found that functionalization of SWCNT by

1‐bromoadamantane leads to formation in inner volume of

nanotube carbon structures with free standing bromine atoms. It

was observed also that during functionalization of the

nanotubes in gas phase by mercury chloride molecules crystal

with parameters of bulk one may be obtained. Loading the

molecules of copper chloride into the nanotubes during gas‐

phase synthesis may lead to crystal lattice transformation from

cubic to rhombohedral symmetry types.


Carbon Materials



4 The structural

peculiarities of nano‐ and

micro‐meter size scale

diamond crystals

Diamond is one of the most well‐known allotropic

modifications of carbon. It is widely applied in industry due to

its exclusive physical properties which include: superior

hardness, biocompatibility, high thermal conductivity and

electrical resistivity, chemical stability, unique properties for

optical elements and coatings [1,39,66,143,144]. There are

various methods for production of diamond, amongst which the

best known is detonation diamonds, and crystals obtained by

CVD [66,67]. Depending on the synthesis methods it is possible

to control size of the producing crystals from nano‐ to micro‐

meter scale. Nanodiamonds, for example, are widely used in

biomedical and mechanical applications [145–147]. Variation of

synthesis conditions allows growth a quasi two‐dimensional

diamonds and single crystal needles with perfect pyramid shape

[148]. Fabrication of well‐shaped and defect‐free diamond single

crystals and especially low‐dimensional diamonds (such as

nanowires and nanoplatelets) is very important for

nanotechnology applications [149,150].

Revealing structural characterization is necessary for

determination of optimal conditions providing fabrication of

materials with desired properties and for correct revealing of

their physical properties. This chapter is focused on description

of study of nanodiamonds structure and its transformations

under heat treatment, as well as on revealing of structural


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peculiarities of quasi two‐dimensional and micro‐meter sized

diamonds crystals.

4.1 THE TRANSFORMATION OF SP3 TO SP2 C-C BOND IN NANODIAMOND UNDER HEAT TREATMENT

An example of the structure transformation of the

nanodiamonds (ND) is formation of so‐called “onion‐like”

structures. The onion‐like carbon (OLC) obtained their name

because of structure resembling concentric layers similar to that

in onion. OLC structures may be produced by graphitization of

nanodiamond and revealed potential for different applications,

including, for example, in energy storage devices, for composite

production and in catalysis [2,151,152]. The sp2/sp3 composites

consisting of one or few nanodiamond cores enveloped with a

few graphene layers and OLC were produced by a controllable

graphitization of explosive NDs in vacuum within the

temperature range of 1200–2140 K [153]. A transformation of the

defective curved graphene sheets into OLC structure occurs

after a complete transformation of diamond cores.

Structure of samples was monitored by Raman spectroscopy

and by high resolution electron microscopy techniques [154–

156]. Nanodiamonds with crystallite size varied from 4 to 12 nm

were analyzed. For all samples, the pristine nanodiamonds were

characterized by two peaks in the Raman spectra. The first one

is a relatively narrow and low intensity Raman band around

1325 cm−1 (D‐mode of sp2‐bonded phase). The second one is a

broad and intense band around 1600 cm−1 (G‐mode of sp2‐

bonded carbon). While the annealing temperature increases D‐

mode shifts towards high frequencies, and G‐mode shifts

towards low frequencies. In addition, after annealing at

temperatures higher than 1800 K a two‐phonon 2D‐mode

appears. The obtained results were interpreted as modifications

occurred at different stages of onion‐like structures formation.

An example of TEM analysis of ND transformation under

heat treatment is presented in Fig. 4‐1. High resolution TEM

Chapter 4: The structural peculiarities of nano‐ and micro‐meter size scale

diamond crystals


image of nanodiamonds before heat treatment procedure is

presented in Fig. 4‐1a. Crystals with average size of 4‐5 nm have

high atomic order which is revealed by lattice fringes of

enlarged HRTEM image and by sharp rings on FFT pattern

corresponding to 111 and 220 interplanar distances in diamond

(Fig. 4‐1a insert). After heat treatment particles possess structure

of graphitic layers rolled into the onion structure (Fig. 4‐1b).

Distance between layers is about 0.33 nm. Small empty hole in

central part of the particles (see enlarged HRTEM) testifies

complete transformation of diamond into graphite‐like structure.

This fact was also confirmed by disappearance of diffraction

peaks in FFT. The schematical representation of ND crystal

structure transformation is illustrated in Fig. 4‐1c. Other details

about nanodiamonds characterizations are presented in

Article VI (see Appendix).

Figure 4‐1. HRTEM images and FFT of nanodiamond before (a) and after (b) heat

treatment. Schematically representation of sp3 to sp2 transformation (c).


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4.2 DIAMOND PLATELETS PRODUCED BY CHEMICAL VAPOR DEPOSITION

Diamond crystal characteristics might be controlled by

varying the CVD parameters such as substrate temperature,

carbon‐containing gas composition and others [66–68]. The most

usual is fabrication of columnar polycrystalline diamond films

consisting of needle‐like crystallites [157,158]. At the same time

production of diamond crystallites with other shapes remains

challenging. This work was dedicated to fabrication of

hexagonal diamond with nanometer thickness and with lateral

size of few micrometers. The material synthesis was performed

using a direct current discharge plasma enhanced chemical

vapor deposition from hydrogen and methane gas mixture. SEM

analysis revealed that obtained diamond films were composed

of several densely packed thin platelets (quasi two‐dimensional

crystal) with lateral size of few micrometers and thickness of

about few tens nanometers (Fig. 4‐2a). TEM analysis of different

diamond platelets revealed absence of significant structural

defects within the crystallites, which indicates their high

structural quality. Characteristics of the registered diffraction

pattern indicate obvious six‐fold rotational symmetry of the

crystallite structure. The interplanar distances in this structure

were evaluated to be 0.1260 nm proving that the platelets are

composed of stacked (111) diamond plane. TEM observations

reveal presence of structural traces caused by planar stacking

faults and twin boundaries in {111} planes (see Fig. 4‐2b, c).

Usage of the HRTEM technique allows analysis the atomic

structure of twinned grain boundary (see Fig. 402d). The

assumption that diamond platelets are composed by stacked

(111) planes is in excellent agreement with previously observed

nanostructures with quasi‐2D shapes for materials with face

centered cubic (FCC) unit cell. On our opinion, the low methane

concentration and relatively high temperature are the key

parameters which provide lateral growth and, as a result,

predominant platelet‐like morphology of the crystallites in CVD

film. Such conditions highly reduce probability of formation of


diamond crystals


three‐atom nucleus on atomically smooth {111} planes. In

combination with absence of twinning on {111} planes (i.e. re‐

entrant corners in {111} area) it causes extremely low growth

rate in <111>directions and strong 111‐faceting of resulting

crystallites.

Figure 4‐2. Typical SEM image (a) and TEM image of cross‐section (b) of the diamond

platelet. (c) Electron diffraction pattern revealed (111) twin’s boundary at cross‐

section of platelet. (d) High resolution bright field TEM image showing the twin

boundary atomic structure. The images reproduced in accord with Article VII (see

in Appendix)

The 2D shapes for FCC nanostructures might be described in

a framework of kinetic and thermodynamic approach of

modified Wulff constructions for twinned nanoparticles.

Accordingly, such type of shapes can be constructed by

considering a particle with several lamellar twins and with fast


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{100} surface growth velocities (i.e. slow {111} growth) and

applying kinetic growth enhancements (at twin plane or/and re‐

entrant surfaces). The resulting structure is also expected to be

hexagonal platelet with (111) main plane.

Thus, in our case formation of the hexagonal diamond

platelets is, apparently, a result of lateral growth between

parallel planar stacking faults in {111} planes in combination

with extremely low growth velocity in <111> directions. The

stacking faults appeared probably at initial stage of the

deposition process. The low growth rate is because of rather

high temperature. More details are presented in Article VII

(see Appendix).

4.3 STRUCTURAL PECULIARITIES OF SINGLE CRYSTAL DIAMOND NEEDLES OF NANOMETER THICKNESS

Fabrication of textured diamond films with controlled

crystallographic orientation of individual crystallites is very

important nowadays because it provides mass production of the

diamond needles with high structural perfection [159]. Such

kind of diamond needles are attractive for different application

including atomic force microscopy, cutting tools etc. [160,161].

The SEM images of diamond needles with shape of perfect

rectangular pyramids are presented in Fig. 4‐3a, b. These

diamond needles were produced by chemical vapor deposition

with appropriate adjustment of the process parameters.

However, some defects may occur during diamond films

fabrication. The structural defects analysis was performed to

understand the relationship between fabrication condition and

diamond structural perfection. Combination of Raman

spectroscopy and transmission electron microscopy techniques

allows revealing and analysis of the crystal defects [82,162].

In this work, we applied some CVD techniques modification

to find new parameters of film growth. During CVD process

nitrogen for a short time was added into gas mixture because it

is known the nitrogen addition changes the ratio of the growth


diamond crystals


rates of {100} and {111} surfaces almost by a factor of about 4.

One may expect that the nitrogen addition could significantly

increase amount of sp2 carbon. It was revealed in our study that

Rama spectra change may be attributed to sp3 type defects

appeared during synthesis without nitrogen addition. The

existence of graphitic carbon (detected by Raman spectroscopy)

may indicate growth conditions which are not enough optimal

for diamond formation. Similar coexistence of diamond (sp3)

and graphitic (sp2) carbon is often observed in CVD diamond

films. Addition of 1% of nitrogen into the gas mixture during

the growth not only leads to an abrupt and substantial decrease

in the diameter of the diamond needle, but also causes changes

in the defect concentration in the growing crystal.


Carbon Materials


Figure 4‐3. SEM images of a diamond with perfect pyramidal crystallite shape (a, b).

HRTEM image reveals a multiple twin’s boundaries/stacking faults propagating in the

crystallites (c). Bright field TEM image of needle apex with variations of the diffraction

contrast, which indicate the presence of numerous extended defects in this area (d).

The images reproduced in accord with Article VIII (see in Appendix)

Transmission electron microscopy provides an information

about atomic structure of the needles. SEM and TEM images

revealed that lateral surface of the diamond needles is generally

rougher in comparison with their basal facets (Fig. 4‐3b, d). This

roughness is mostly determined by the oxidation step of the

fabrication process. Thorough oxidation results in almost flat

surface; while insufficient oxidation cycle leaves residuals of the


diamond crystals


dendrite‐like shape crystallites. The needle itself is represented

by a perfect undisturbed diamond crystal, what is evidenced by

the large area convergent beam electron diffraction (LACBED)

pattern with straight and sharp high‐order Laue‐zone

diffraction. According to LACBED the pyramid base is formed

by {100} plane, and the lateral surfaces are close to {110} planes.

The observed (001) texture of obtained films indicates that

applied conditions of CVD process provide the slowest crystal

growth along ⟨100⟩ crystallographic direction. The fastest growing direction at these conditions is ⟨110⟩ leading to degeneration of {110} facets into edges between {100} and {111}

facets.

The dendrites structure is dramatically differ from the rest of

the needle in crystal quality—they acquire a high density of

stacking faults and {111} twins, thus creating a high

concentration of planar defects in the near surface region of the

needles (Fig. 4‐3c). The distance between stacking faults/twin

boundaries is averaged at about 5 nm. Figure 4‐3d shows the

region of the needle near the apex. The crystal here contains

much more defects, which are visible as the irregularity of

thickness fringes.

Raman scattering analysis reveals a few peculiar features

related to the defects found in the needle. Incorporation of

possible point defects into the volume of growing crystal may be

explained on the basis of detailed molecular modeling of the

diamond growth on {100} surface. By kinetic analysis it may be

shown, that migration of CH2 fragments leads to their

alignment in dimers chain direction on the reconstructed

C{100}:H 2 × 1 surface. Aligned CH2 bridges inevitably contain

one atom gaps. Thus, even at moderate growth rates, there is a

high probability that these monoatomic voids are trapped inside

the crystal, forming vacancy type point defects. High growth

rate upon nitrogen addition may be also the reason for sp2

fragments inclusion into bulk of the diamond crystal showing

broad Raman lines in 1500–1600 cm−1 spectral range. Further

details about this structural study is presented in Article VIII

(see Appendix).


Carbon Materials



In this Chapter, it was briefly shown that applying various

synthesis methods diamonds crystals from nano‐ to micro‐mete

crystal sizes could be grown.

It was shown that the heat treatment of nanodiamonds leads

to gradually transformation of carbon‐to‐carbon bonds sp3 to sp2,

and thus, to carbon onion‐like structure formation.

Adjustment of CVD process parameters allows formation of

quasi two‐dimensional hexagonal diamond platelets with

several micrometers in size. TEM structure analysis revealed

that platelets are consists of stacked (111) atomic planes.

It was demonstrated that needle grown by optimized CVD

possess a perfect undisturbed diamond crystal structure. The

needles have rough surface which consist of high density of

stacking faults and {111} twins.


5 Conclusions

In this Thesis, experimental analysis of structural

peculiarities of various carbon materials were presented. The

comprehensive analysis of electron microscopy methodology

was performed. Corresponding adjustment of spherical

aberration corrector allowed improvement of TEM instrument

parameters with increase of signal‐to‐noise ratio up to 30%. This

TEM tuning allowed revealing of important peculiarities in

structures of carbon materials, including:

‐ formation of carbon structures with free‐standing isolated

Br atoms in 1‐bromoadamantane functionalized single

walled carbon nanotubes;

‐ formation of ordered compounds of guest molecules of

different types inside single wall carbon nanotubes and

detection their structural transformation during synthesis;

‐ gradual transformation during heat treatment sp3

carbon‐to‐carbon bonds of nanodiamond particles into

sp2 and formation onion‐like carbon structures formation;

‐ obtaining quasi two‐dimensional hexagonal diamond

platelets of micrometer scale by chemical vapor

deposition and revealing their structure consisting of

stacked (111) atomic planes;

‐ determination of structural characteristics of single‐

crystal diamond needles and elicitation of high density

stacking faults and {111} twins in these diamond

crystallites.


Carbon Materials



References

1. Pierson H.O. Handbook of Carbon, Graphite, Diamond

and Fullerenes // Handb. Carbon, Graph. Diam. Fullerenes. 1993.

P. 25–69.

2. Rao C.N.R. et al. Fullerenes, nanotubes, onions and

related carbon structures // Materials Science and Engineering R.

1995. Vol. 15, № 6. P. 209–262.

3. Heimann R.B., Evsvukov S.E., Koga Y. Carbon allotropes:

a suggested classification scheme based on valence orbital

hybridization // Carbon N. Y. 1997. Vol. 35, № 10–11. P. 1654–

1658.

4. Suarez‐Martinez I., Grobert N., Ewels C.P. Nomenclature

of sp 2 carbon nanoforms // Carbon. 2012. Vol. 50, № 3. P. 741–

747.

5. Peschel G. Carbon‐carbon bonds Hybridization // Freie

Univ. Berlin website. 2005.

6. Smith P.P., Buseck P.R. Graphitic carbon in the Allende

meteorite: a microstructural study. // Science (80‐. ). 1981. Vol.

212, № 4492. P. 322–324.

7. Smith P.P., Buseck P.R. Carbyne forms of carbon: do they

exist? // Science. 1982. Vol. 216, № 4549. P. 984–986.

8. Novoselov K.S. et al. Electric field effect in atomically

thin carbon films. // Science. 2004. Vol. 306, № 5696. P. 666–669.

9. Geim A.K., Novoselov K.S. The rise of graphene // Nat.

Mater. 2007. Vol. 6, № 3. P. 183–191.

10. Eizenberg M., Blakely J.M. Carbon monolayer phase

condensation on Ni(111) // Surf. Sci. 1979. Vol. 82, № 1. P. 228–

236.

11. Isett L.C., Blakely J.M. Segregation isosteres for carbon at

the (100) surface of nickel // Surf. Sci. 1976. Vol. 58, № 2. P. 397–

414.

12. J. C. Meyer et al. The structure of suspended graphene

sheets // Nature. 2007. Vol. 446, № 1. P. 60–63.


Carbon Materials


13. Carlsson J.M. Graphene: Buckle or break // Nat. Mater.

2007. Vol. 6, № 11. P. 801–802.

14. Geim A.K., Novoselov K.S. The rise of graphene // Nat.

Mater. 2007. Vol. 6, № 3. P. 183–191.

15. Fasolino a, Los J.H., Katsnelson M.I. Intrinsic ripples in

graphene // Nat. Mater. 2007. Vol. 6, № 11. P. 858–861.

16. Wyckoff R.W.G. Crystal Structures. 2nd ed. New York:

Interscience, 1964.

17. Gray D., Mccaughan A., Mookerji B. Crystal Structure of

Graphite , Graphene and Silicon // Phys. Solid State Appl. 2009.

Vol. 2. P. 3–5.

18. International Tables for Crystallography. 2nd Editio / ed.

Wondratschek H., Muller U. New York: Wiley, 2006. 6078 p.

19. Kopelevich Y., Esquinazi P. Graphene Physics in

Graphite // Adv. Mater. 2007. Vol. 19. P. 9.

20. Chung D.D.L. Review: Graphite // Journal of Materials

Science. 2002. Vol. 37, № 8. P. 1475–1489.

21. Popov V.N. Carbon nanotubes: Properties and

application // Materials Science and Engineering R: Reports.

2004. Vol. 43, № 3. P. 61–102.

22. Monthioux M. Introduction to Carbon Nanotubes //

Carbon Meta‐Nanotubes: Synthesis, Properties and Applications.

2011. P. 7–39.

23. Sinclair J. An Introduction to Carbon Nanotubes //

Carbon Nanotubes. 2009. P. 1–8.

24. Murakami Y. et al. Direct synthesis of high‐quality

single‐walled carbon nanotubes on silicon and quartz substrates

// Chem. Phys. Lett. 2003. Vol. 377, № 1–2. P. 49–54.

25. Thostenson E.T., Ren Z., Chou T.‐W. Advances in the

science and technology of carbon nanotubes and their

composites: a review // Compos. Sci. Technol. 2001. Vol. 61, №

13. P. 1899–1912.

26. Sun Y.P. et al. Functionalized carbon nanotubes:

Properties and applications // Acc. Chem. Res. 2002. Vol. 35, №

12. P. 1096–1104.

27. Darwish A.D. Fullerenes // Annu. Reports Sect. “A”

(Inorganic Chem. 2012. Vol. 108. P. 464.

References


28. Birkett P.R. Fullerenes // Annu. Reports Sect. “A”

(Inorganic Chem. 2006. Vol. 102. P. 420.

29. Howard J.B. et al. Fullerenes C60 and C70 in flames. //

Nature. 1991. Vol. 352, № 6331. P. 139–141.

30. Li C.‐Z., Yip H.‐L., Jen A.K.‐Y. Functional fullerenes for

organic photovoltaics // J. Mater. Chem. 2012. Vol. 22, № 10. P.

4161.

31. Chuvilin A. et al. Direct transformation of graphene to

fullerene // Nat. Chem. 2010. Vol. 2, № 6. P. 450–453.

32. Banhart F., Kotakoski J., Krasheninnikov A. V. Structural

defects in graphene // ACS Nano. 2011. Vol. 5, № 1. P. 26–41.

33. Schniepp H.C. et al. Functionalized single graphene

sheets derived from splitting graphite oxide // J. Phys. Chem. B.

2006. Vol. 110, № 17. P. 8535–8539.

34. Ramanathan T. et al. Functionalized graphene sheets for

polymer nanocomposites // Nat. Nanotechnol. 2008. Vol. 3, № 6.

P. 327–331.

35. Wei W., Qu X. Extraordinary physical properties of

functionalized graphene // Small. 2012. Vol. 8, № 14. P. 2138–

2151.

36. Song M., Cai D. Graphene Functionalization: A Review //

Polymer‐ Graphene Nanocomposites. 2012. Vol. 26, № 26. 1‐51 p.

37. Burkhard G. et al. Formation of Cubic Carbon by

Dynamic Shock Compression of a Diamond/Amorphous Carbon

Powder Mixture // Jpn. J. Appl. Phys. 1994. Vol. 33, № 10. P.

5875–5885.

38. Edwards D.F., Philipp H.R. Cubic carbon (diamond) //

Handbook of Optical Constants of Solids. 2012. Vol. 1. P. 665–

673.

39. Bundy F.P. Hexagonal Diamond—A New Form of

Carbon // J. Chem. Phys. 1967. Vol. 46, № 9. P. 3437.

40. Chiem C. et al. Lonsdaleite diamond growth on

reconstructed Si (100) by hot‐filament chemical vapor deposition

(HFCVD) // Korean J. Chem. Eng. 2003. Vol. 20, № 6. P. 1154–

1157.


Carbon Materials


41. Brown G. et al. Electron beam induced in situ

clusterisation of 1D ZrCl4 chains within single‐walled carbon

nanotubes // Chem. Commun. 2001. № 9. P. 845–846.

42. Sceats E.L., Green J.C., Reich S. Theoretical study of the

molecular and electronic structure of one‐dimensional crystals

of potassium iodide and composites formed upon intercalation

in single‐walled carbon nanotubes // Phys. Rev. B ‐ Condens.

Matter Mater. Phys. 2006. Vol. 73, № 12.

43. Ajayan P.M., Iijima S. Capillarity‐induced filling of

carbon nanotubes // Nature. 1993. Vol. 361, № 6410. P. 333–334.

44. Chaturvedi P. et al. Carbon nanotube ‐ Purification and

sorting protocols // Def. Sci. J. 2008. Vol. 58, № 5. P. 591–599.

45. Rahman M.M. et al. Electric and magnetic properties of

Co‐filled carbon nanotube // J. Phys. Soc. Japan. 2005. Vol. 74, №

2. P. 742–745.

46. Sloan J. et al. The size distribution, imaging and

obstructing properties of C60 and higher fullerenes formed

within arc‐grown single walled carbon nanotubes // Chem. Phys.

Lett. 2000. Vol. 316, № January. P. 191.

47. Eliseev A. a et al. Preparation and properties of single‐

walled nanotubes filled with inorganic compounds // Russ.

Chem. Rev. 2009. Vol. 78, № 9. P. 833–854.

48. Hall T. Ultrahigh‐Pressure Research // Science (80‐. ).

1958. Vol. 128, № 3322. P. 445–449.

49. Merlen A. et al. High pressure‐high temperature

synthesis of diamond from single‐wall pristine and iodine

doped carbon nanotube bundles // Carbon N. Y. 2009. Vol. 47,

№ 7. P. 1643–1651.

50. Kawasaki S. et al. Hardness of high‐pressure high‐

temperature treated single‐walled carbon nanotubes // Phys. B

Condens. Matter. 2007. Vol. 388, № 1–2. P. 59–62.

51. Epanchintsev O.G. et al. Highly‐efficient shock‐wave

diamond synthesis from fullerenes // J. Phys. Chem. Solids. 1997.

Vol. 58, № 11. P. 1785–1788.

52. Kenkmann T., Hornemann U., Stoffler D. Experimental

shock synthesis of diamonds in a graphite gneiss // Meteorit.

Planet. Sci. 2005. Vol. 40, № 9–10. P. 1299–1310.

References


53. Greiner N.R. et al. Diamonds in detonation soot // Nature.

1988. Vol. 333, № 6172. P. 440–442.

54. Danilenko V. V. Specific features of synthesis of

detonation nanodiamonds // Combust. Explos. Shock Waves.

2005. Vol. 41, № 5. P. 577–588.

55. Churilov G.N. Synthesis of Fullerenes and Other

Nanomaterials in Arc Discharge // Fullerenes, Nanotub. Carbon

Nanostructures. 2008. Vol. 16, № 5–6. P. 395–403.

56. Subrahmanyam K.S. et al. Simple method of preparing

graphene flakes by an arc‐discharge method // J. Phys. Chem. C.

2009. Vol. 113, № 11. P. 4257–4259.

57. Ando Y., Zhao X. Synthesis of Carbon Nanotubes by

Arc‐Discharge Method // New Diam. Forntier Carbon Technol.

2006. Vol. 16, № 3. P. 123–137.

58. Antisari M.V., Marazzi R., Krsmanovic R. Synthesis of

multiwall carbon nanotubes by electric arc discharge in liquid

environments // Carbon N. Y. 2003. Vol. 41, № 12. P. 2393–2401.

59. Arora N., Sharma N.N. Arc discharge synthesis of carbon

nanotubes: Comprehensive review // Diamond and Related

Materials. 2014. Vol. 50. P. 135–150.

60. Amans D. et al. Nanodiamond synthesis by pulsed laser

ablation in liquids // Diam. Relat. Mater. 2009. Vol. 18, № 2–3. P.

177–180.

61. Thongpool V., Asanithi P., Limsuwan P. Synthesis of

carbon particles using laser ablation in ethanol // Procedia

Engineering. 2012. Vol. 32. P. 1054–1060.

62. Shibagaki K. et al. Synthesis of heavy carbon clusters by

laser ablation in vacuum // Japanese J. Appl. Physics, Part 2 Lett.

2001. Vol. 40, № 8 B.

63. Chen C., Chen W., Zhang Y. Synthesis of carbon nano‐

tubes by pulsed laser ablation at normal pressure in metal nano‐

sol // Phys. E Low‐Dimensional Syst. Nanostructures. 2005. Vol.

28, № 2. P. 121–127.

64. HORNBOSTEL B. et al. Arc Discharge and Laser

Ablation Synthesis of Singlewalled Carbon Nanotubes // Carbon

Nanotub. 2006. Vol. 1. P. 1–18.


Carbon Materials


65. Carvalho A.F. et al. Simultaneous CVD synthesis of

graphene‐diamond hybrid films // Carbon N. Y. 2016. Vol. 98. P.

99–105.

66. Gicquel A. et al. CVD diamond films: From growth to

applications // Curr. Appl. Phys. 2001. Vol. 1, № 6. P. 479–496.

67. Schwarz S. et al. CVD‐diamond single‐crystal growth // J.

Cryst. Growth. 2004. Vol. 271, № 3–4. P. 425–434.

68. Gracio J.J., Fan Q.H., Madaleno J.C. Diamond growth by

chemical vapour deposition // J. Phys. D. Appl. Phys. 2010. Vol.

43, № 37. P. 374017.

69. Muñoz R., Gómez‐Aleixandre C. Review of CVD

synthesis of graphene // Chemical Vapor Deposition. 2013. Vol.

19, № 10–12. P. 297–322.

70. Zheng B., Li Y., Liu J. CVD synthesis and purification of

single‐walled carbon nanotubes on aerogel‐supported catalyst //

Appl. Phys. A Mater. Sci. Process. 2002. Vol. 74, № 3. P. 345–348.

71. Kuwana K., Saito K. Modeling CVD synthesis of carbon

nanotubes: Nanoparticle formation from ferrocene // Carbon N.

Y. 2005. Vol. 43, № 10. P. 2088–2095.

72. Hsin Y.L. et al. Production and in‐situ metal filling of

carbon nanotubes in water // Adv. Mater. 2001. Vol. 13, № 11. P.

830–833.

73. Gately R.D., In het Panhuis M. Filling of carbon

nanotubes and nanofibres // Beilstein Journal of

Nanotechnology. 2015. Vol. 6, № 1. P. 508–516.

74. Tsang S.C. et al. A simple chemical method of opening

and filling carbon nanotubes // Nature. 1994. Vol. 372, № 6502. P.

159–162.

75. Chancolon J. et al. Filling of carbon nanotubes with

selenium by vapor phase process. // J. Nanosci. Nanotechnol.

2006. Vol. 6, № 1. P. 82–86.

76. Fedotov P. V. et al. Optical properties of single‐walled

carbon nanotubes filled with CuCl by gas‐phase technique //

Phys. Status Solidi Basic Res. 2014. Vol. 251, № 12. P. 2466–2470.

77. Monthioux M., Flahaut E., Cleuziou J.‐P. Hybrid carbon

nanotubes: Strategy, progress, and perspectives // J. Mater. Res.

2006. Vol. 21, № 11. P. 2774–2793.

References


78. Sloan J. et al. Structural changes induced in nanocrystals

of binary compounds confined within single walled carbon

nanotubes: A brief review // Inorganica Chimica Acta. 2002. Vol.

330, № 1. P. 1–12.

79. Kuzmany H. et al. Raman spectroscopy of fullerenes and

fullerene‐nanotube composites. // Philos. Trans. A. Math. Phys.

Eng. Sci. 2004. Vol. 362, № 1824. P. 2375–2406.

80. Dresselhaus, M.S., Dresselhaus, G., Eklund P.C. Raman

Scattering in Fullerenes // J. Phys. Chem. 1996. Vol. 27, №

October 1995. P. 351–371.

81. Dresselhaus M.S. et al. Perspectives on carbon nanotubes

and graphene Raman spectroscopy // Nano Letters. 2010. Vol. 10,

№ 3. P. 751–758.

82. Pimenta M.A. et al. Studying disorder in graphite‐based

systems by Raman spectroscopy // Phys. Chem. Chem. Phys.

2007. Vol. 9, № 11. P. 1276–1291.

83. Li Z.Q. et al. X‐ray diffraction patterns of graphite and

turbostratic carbon // Carbon N. Y. 2007. Vol. 45, № 8. P. 1686–

1695.

84. Reznik D. et al. X‐ray powder diffraction from carbon

nanotubes and nanoparticles // Phys. Rev. B. 1995. Vol. 52, № 1.

P. 116–124.

85. Tomita S. et al. Diamond nanoparticles to carbon onions

transformation: X‐ray diffraction studies // Carbon N. Y. 2002.

Vol. 40, № 9. P. 1469–1474.

86. Cuesta A. et al. Comparative performance of X‐ray

diffraction and Raman microprobe techniques for the study of

carbon materials // J. Mater. Chem. 1998. Vol. 8. P. 2875–2879.

87. Van Tendeloo G., Amelinckx S. Electron Microscopy of

Fullerenes and Related Materials // Characterization of

Nanophase Materials. 2001. Vol. 1. 353‐393 p.

88. Magni S. et al. FIB/SEM characterization of carbon‐based

fibers // Scanning. 2007. Vol. 29, № 4. P. 185–195.

89. Wunderlich W., Foitzik A.H., Heuer A.H. On the

quantitative EDS analysis of low carbon concentrations in

analytical TEM // Ultramicroscopy. 1993. Vol. 49, № 1–4. P. 220–

224.


Carbon Materials


90. University of California Riverside. Introduction to

Energy Dispersive X‐ray Spectrometry (EDS) // Cent. Facil. Adv.

Microsc. Microanal. 2011. P. 1–12.

91. Newbury D.E., Ritchie N.W.M. Elemental mapping of

microstructures by scanning electron microscopy‐energy

dispersive X‐ray spectrometry (SEM‐EDS): extraordinary

advances with the silicon drift detector (SDD) // J. Anal. At.

Spectrom. 2013. Vol. 28, № 7. P. 973–988.

92. Newbury D.E., Ritchie N.W.M. Is scanning electron

microscopy/energy dispersive X‐ray spectrometry (SEM/EDS)

quantitative? // Scanning. 2013. Vol. 35, № 3. P. 141–168.

93. Yoshizawa N. et al. TEM and electron tomography

studies of carbon nanospheres for lithium secondary batteries //

Carbon N. Y. 2006. Vol. 44, № 12. P. 2558–2564.

94. Meyer J.C. Transmission electron microscopy (TEM) of

graphene // Graphene. 2014. P. 101–123.

95. Sinclair R., Itoh T., Chin R. In Situ TEM Studies of Metal‐

Carbon Reactions // Microscopy and Microanalysis. 2002. Vol. 8,

№ 4. P. 288–304.

96. Safarova K., Drovak a, Kubinek R. Usage of AFM, SEM

and TEM for the research of carbon nanotubes // Mod. Res. Educ.

Top. Microsc. 2007. P. 513–519.

97. Zhang B., Su D.S. Transmission electron microscopy and

the science of carbon nanomaterials // Small. 2014. Vol. 10, № 2.

P. 222–229.

98. Yehliu K., Vander Wal R.L., Boehman A.L. Development

of an HRTEM image analysis method to quantify carbon

nanostructure // Combust. Flame. 2011. Vol. 158, № 9. P. 1837–

1851.

99. Vander Wal R.L. et al. Analysis of HRTEM images for

carbon nanostructure quantification // J. Nanoparticle Res. 2004.

Vol. 6, № 6. P. 555–568.

100. Kuwahara S., Sugai T., Shinohara H. A new AFM‐

HRTEM combined technique for probing isolated carbon

nanotubes. // Nanotechnology. 2009. Vol. 20, № 22. P. 225702.

101. Williams D., Carter C. The transmission electron

microscope. 1996.

References


102. Pfaff M. et al. Low‐energy electron scattering in carbon‐

based materials analyzed by scanning transmission electron

microscopy and its application to sample thickness

determination // J. Microsc. 2011. Vol. 243, № 1. P. 31–39.

103. Song F.Q. et al. Free‐standing graphene by scanning

transmission electron microscopy // Ultramicroscopy. 2010. Vol.

110, № 12. P. 1460–1464.

104. Sourty E. et al. High‐angle annular dark field scanning

transmission electron microscopy on carbon‐based functional

polymer systems. // Microsc. Microanal. 2009. Vol. 15, № 3. P.

251–258.

105. Bachtold a et al. Scanned probe microscopy of electronic

transport in carbon nanotubes. // Phys. Rev. Lett. 2000. Vol. 84,

№ 26. P. 6082–6085.

106. de Broglie L. A tentative theory of light quanta // Phil.

Mag. 1924. Vol. 47. P. 446–458.

107. Bragg W. The diffraction of short electromagnetic waves

by a crystal // Proc. Camb. Philol. Soc. 1913. Vol. 17. P. 43–57.

108. Stadelmann P. JEMS electron microscopy simulation

software. 2004.

109. Scherzer O. Uber einige Fehler von Elektronenlinsen //

Zeitschrift fur Phys. 1936. Vol. 101, № 9–10. P. 593–603.

110. Brandon D., Kaplan W.D. Microstructural

Characterization of Materials // Microstructural Characterization

of Materials. 2nd Editio. England: John Wiley & Sons Ltd, 2008.

536 p.

111. Shindo D., Oikawa T. Analytical Electron Microscopy for

Materials Science. Japan: Springer Japan, 2002. 161 p.

112. Scherzer O. Spharische und chromatische Korrektur von

Elektronenlinsen // Optik (Stuttg). 1947. Vol. 2. P. 114–132.

113. Rose H. Outline of a spherically corrected semiaplanatic

medium‐voltage transmission electron microscope // Optik

(Stuttg). 1990. Vol. 85. P. 19–24.

114. Krivanek O.L. et al. Aberration correction in the STEM //

Inst. Phys. Conf. Ser. EMAG97. 1997. Vol. 153, № Section 2. P.

35–40.


Carbon Materials


115. Thon F. Zur Defokussierungsabhängigkeit des

Phasenkontrastes bei der elektronmikroskopischen Abbildung //

Z. Naturforsch. 1966. Vol. 21a. P. 476–478.

116. Krivanek O.L. A method for determining the coefficient

of spherical aberration from a single electron micrograph //

Optik (Stuttg). 1976. Vol. 45, № 1. P. 97–101.

117. Zemlin F. et al. Coma‐free alignment of high resolution

electron microscopes with the aid of optical diffractograms //

Ultramicroscopy. 1978. Vol. 3, № C. P. 49–60.

118. Sidorov M.V. CtfExplorer software. 2002.

119. Krivanek O.L., Mooney P.E. Applications of slow‐scan

CCD cameras in transmission electron microscopy //

Ultramicroscopy. 1993. Vol. 49, № 1–4. P. 95–108.

120. Stenkamp D. Detection and quantitative assessment of

image aberrations from single HRTEM lattice images // Journal

of Microscopy. 1998. Vol. 190, № 1–2. P. 194–203.

121. Uhlemann S., Haider M. Residual wave aberrations in

the first spherical aberration corrected transmission electron

microscope // Ultramicroscopy. 1998. Vol. 72, № 3–4. P. 109–119.

122. Erni R. Aberration‐Corrected Imaging in Transmission

Electron Microscopy. London: Imperial College Press, 2010. 335

p.

123. Lee Z. et al. Electron dose dependence of signal‐to‐noise

ratio, atom contrast and resolution in transmission electron

microscope images // Ultramicroscopy. 2014. Vol. 145. P. 3–12.

124. Rose H. Theoretical aspects of image formation in the

aberration‐corrected electron microscope // Ultramicroscopy.

2010. Vol. 110, № 5. P. 488–499.

125. Haider M. et al. Information transfer in a TEM corrected

for spherical and chromatic aberration. // Microsc. Microanal.

2010. Vol. 16. P. 393–408.

126. Haider M., Uhlemann S., Zach J. Upper limits for the

residual aberrations of a high‐resolution aberration‐corrected

STEM // Ultramicroscopy. 2000. Vol. 81, № 3–4. P. 163–175.

127. Chang L.Y., Kirkland A.I., Titchmarsh J.M. On the

importance of fifth‐order spherical aberration for a fully

References


corrected electron microscope // Ultramicroscopy. 2006. Vol. 106,

№ 4–5. P. 301–306.

128. Lupini A.R., Pennycook S.J. Tuning Fifth‐Order

Aberrations in a Quadrupole‐Octupole Corrector // Microsc.

Microanal. 2012. Vol. 18, № 4. P. 699–704.

129. Benard Y., Lopez‐Gil N., Legras R. Subjective depth of

field in presence of 4th‐order and 6th‐order Zernike spherical

aberration using adaptive optics technology // J. Cataract Refract.

Surg. 2010. Vol. 36, № 12. P. 2129–2138.

130. Urban K.W. et al. Negative spherical aberration

ultrahigh‐resolution imaging in corrected transmission electron

microscopy. // Philos. Trans. A. Math. Phys. Eng. Sci. 2009. Vol.

367, № 1903. P. 3735–3753.

131. Jia C.L. et al. On the benefit of the negative‐spherical‐

aberration imaging technique for quantitative HRTEM //

Ultramicroscopy. 2010. Vol. 110, № 5. P. 500–505.

132. Wei D. et al. Ultrathin rechargeable all‐solid‐state

batteries based on monolayer graphene // J. Mater. Chem. A.

2013. Vol. 1, № 9. P. 3177–3181.

133. Barthel J., Thust A. Quantification of the information

limit of transmission electron microscopes // Phys. Rev. Lett.

2008. Vol. 101, № 20.

134. Lee Z. et al. Optimum HRTEM image contrast at 20 kV

and 80 kV‐‐exemplified by graphene. // Ultramicroscopy. 2012.

Vol. 112, № 1. P. 39–46.

135. Tonkikh A.A. et al. Metallization of single‐wall carbon

nanotube thin films induced by gas phase iodination // Carbon

N. Y. 2015. Vol. 94. P. 768–774.

136. Janas D. et al. Iodine monochloride as a powerful

enhancer of electrical conductivity of carbon nanotube wires //

Carbon N. Y. 2014. Vol. 73. P. 225–233.

137. Tonkikh A.A. et al. Optical spectroscopy of iodine‐doped

single‐wall carbon nanotubes of different diameter // Phys.

Status Solidi Basic Res. 2012. Vol. 249, № 12. P. 2454–2459.

138. Bendiab N. et al. Structural determination of iodine

localization in single‐walled carbon nanotube bundles by


Carbon Materials


diffraction methods // Phys. Rev. B. 2004. Vol. 69, № 19. P.

195415.

139. Guan L. et al. Polymorphic structures of iodine and their

phase transition in confined nanospace // Nano Lett. 2007. Vol. 7,

№ 6. P. 1532–1535.

140. Alvarez L. et al. High‐pressure behavior of polyiodides

confined into single‐walled carbon nanotubes: A Raman study //

Phys. Rev. B ‐ Condens. Matter Mater. Phys. 2010. Vol. 82, № 20.

141. Kotaka Y., Yamazaki T., Kataoka Y. Atomic‐resolution

imaging and analysis with cs‐corrected scanning transmission

electron microscopy // Fujitsu Sci. Tech. J. 2010. Vol. 46, № 3. P.

249–256.

142. Wilson R.H., Kasper J.S. Golden Book of Phase

Transitions. Wroclaw, 2002. 123 p.

143. Zaitsev A.M. Optical Properties of Diamond // Book.

2001. P. 502.

144. Robertson J. Properties of diamond‐like carbon // Surface

and Coatings Technology. 1992. Vol. 50, № 3. P. 185–203.

145. Kong X., Cheng P. Application of Nanodiamonds in

Biomolecular Mass Spectrometry // Materials (Basel). 2010. Vol.

3, № 3. P. 1845–1862.

146. Passeri D. et al. Biomedical Applications of

Nanodiamonds: An Overview // J. Nanosci. Nanotechnol. 2015.

Vol. 15, № 2. P. 972–988.

147. Dolmatov V.Y. Detonation‐synthesis nanodiamonds:

synthesis, structure, properties and applications // Russ. Chem.

Rev. 2007. Vol. 76, № 4. P. 339–360.

148. Obraztsov A.N. et al. Production of single crystal

diamond needles by a combination of CVD growth and thermal

oxidation // Diam. Relat. Mater. 2009. Vol. 18, № 10. P. 1289–

1293.

149. Aharonovich I. et al. Bottom‐up engineering of diamond

micro‐ and nano‐structures // Laser Photon. Rev. 2013. Vol. 7, №

5. P. L61–L65.

150. Aharonovich I., Neu E. Diamond nanophotonics // Adv.

Opt. Mater. 2014. Vol. 2. P. 911–928.

References


151. Su D. et al. Oxidative dehydrogenation of ethylbenzene

to styrene over ultra‐dispersed diamond and onion‐like carbon

// Carbon N. Y. 2007. Vol. 45, № 11. P. 2145–2151.

152. Pech D. et al. Ultrahigh‐power micrometre‐sized

supercapacitors based on onion‐like carbon // Nat. Nanotechnol.

2010. Vol. 5. P. 651–654.

153. Kuznetsov V. et al. Controllable electromagnetic

response of onion‐like carbon based materials // Physica Status

Solidi (B) Basic Research. 2008. Vol. 245, № 10. P. 2051–2054.

154. Rao, A.M.; Richter, E.; Bandow, S.; Chase, B.; Eklund,

P.C.; Williams, K.A.; Fang, S.; Subbaswamy, K.R.; Menon, M.;

Thess, A.; Smalley, R.E.; Dresselhaus, G.; Dresselhaus M.S.

Diameter‐selective Raman scattering from vibrational modes in

carbon nanotubes // Science (80‐. ). 1997. Vol. 275, № January. P.

187–191.

155. Rao A.M. et al. Evidence for charge transfer in doped

carbon nanotube bundles from Raman scattering // Nature. 1997.

Vol. 388, № 6639. P. 257–259.

156. Pfeiffer R. et al. Evidence for trans‐polyacetylene in

nano‐crystalline diamond films // Diam. Relat. Mater. 2003. Vol.

12, № 3–7. P. 268–271.

157. Zolotukhin A. et al. Thermal oxidation of CVD diamond

// Diam. Relat. Mater. 2010. Vol. 19, № 7–9. P. 1007–1011.

158. Zolotukhin A.A. et al. Single‐crystal diamond

microneedles shaped at growth stage // Diam. Relat. Mater. 2014.

Vol. 42. P. 15–20.

159. Ding M.Q., Li L., Feng J. A study of high‐quality

freestanding diamond films grown by MPCVD // Appl. Surf. Sci.

2012. Vol. 258, № 16. P. 5987–5991.

160. Obraztsov A.N. et al. Single crystal diamond tips for

scanning probe microscopy. // Rev. Sci. Instrum. 2010. Vol. 81,

№ 1. P. 13703.

161. Tuyakova F.T., Obraztsova E.A., Ismagilov R.R. Single‐

crystal diamond pyramids: synthesis and application for atomic

force microscopy // J. Nanophot. 2016. Vol. 10, № 1. P. 12517.


Carbon Materials


162. Knight D.S., White W.B. Characterization of diamond

films by Raman spectroscopy // J. Mater. Res. 1989. Vol. 4, № 2.

P. 385–393.

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ISBN 978-952-61-2381-3ISSN 1798-5668


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ANDREY OREKHOV

ELECTRON MICROSCOPY STUDY OF STRUCTURALPECULIARITIES OF CARBON MATERIALS


This work presents results of high resolution transmission electron microscopy analysis

of structural peculiarities of some of nanocarbons materials forms. It was

performed the optimization of optical system of aberration corrected transmission electron

microscope to increase of the signal-to-noise ratio. The optimized TEM instruments

were used in this work for structural characterization of the nanodiamonds, two-

dimensional (2D) structures and needle-like diamonds, onion-like nanocarbons and other graphene-based structures, including

composites consisting of linear (1D) CuCl and Hg2Cl2 crystals encapsulated in single-walled carbon nanotubes. Obtained results allowed

appropriate development of production processes for these carbon nanostructures and


ANDREY OREKHOV

dissertations in forestry and natural sciences

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