Fabrication and properties of diamond-like carbon films

in discharge plasmas

Maksym Rybachuk

A thesis submitted for the degree of Doctor of Philosophy

at the Queensland University of Technology

Australia

March 2007

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Declaration of originality

The work presented in this thesis is the original work of the author carried out under

the guidance of Prof. John M. Bell. Tasks related to fabrication, measurements and

analysis of the films were performed solely by the author. The work contained in this

thesis has not been previously submitted to meet requirements for an award at this or

any other higher education institution. To the best of my knowledge and belief, the

thesis contains no material previously published or written by another person except

where due reference is made.

Maksym Rybachuk

Date: 21 March 2007

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Acknowledgements

I would like to thank and acknowledge the help and support of the following people:

• My principal supervisor John Bell and my associate supervisor John

Kavanagh for friendship, guidance, mentorship and for providing invaluable

support during all stages of the project.

• I am very thankful to all staff members of Laserdyne Pty Ltd and Russ Helms

for kindly providing support during the large part of the experimental work.

• To colleagues and friends at the School of Chemistry at QUT and in

particular to Llew Rintoul for the assistance with IR and Raman

measurements, to Geoff Will, Peter Frederics and Eric Waclawick for

reviewing this thesis and to Thor Bostrom for the assistance with the SEM

measurements.

• To Barry Wood of the University of Queensland for the assistance with the

XPS measurements

• To John Drennan (Nano-MNRF) for the financial support provided via TAP

program.

• To Steven Prawer and Alberto Cimmino of the University of Melbourne for

the assistance with 244 nm UV Raman measurements.

• To Greg Hope of the Griffith University for the assistance with 325 nm UV

Raman measurements.

• To Nunzio Motta for the assistance with STM, STS and AFM.

• To Federico Rosei of the Institut National de la Recherche Scientifique,

University of Québec for guidance.

• To Richard Brown for encouraging to embark on a PhD road.

• To fellow postgraduate students: Roland Goh, Cameron Brown, Nick

Gaddum, Lorne Gale, Praveen Posinasetti, Nick Ward, Jay Madhani and to

many others for their friendship and support.

• To my parents Dmitriy and Vera Rybachuk for their absolute support.

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___________________________________________________Author’s publications

Refereed journal publications

12 M. Rybachuk, J.M. Bell "Polyacetylene and poly(p-phenylene vinylene) π-

conjugated chains and sp- bonded chains probed by Resonant Raman scattering in

diamond-like carbon" Carbon, submitted 6/2008

11 M. Rybachuk, A. Hu, J.M. Bell "Resonant Raman scattering from

polyacetylene and poly(p-phenylene vinylene) chains intercalated into hydrogenated

diamond-like carbon" Applied Physics Letters, submitted 3/2008

10 A. Hu, M. Rybachuk, Q. - B. Lu, W. W. Duley “Femtosecond pulsed laser

deposition and optical properties of diamond-like amorphous carbon films embedded

with sp-bonded carbon chains”, Diamond and Related Materials, accepted for

publication 3/2008, doi: 10.1016/j.diamond.2008.03.024

9 A. Hu, M. Rybachuk, Q.-B. Lu, and W. W. Duley, "Direct synthesis of sp-

bonded carbon chains on graphite surface by femtosecond laser irradiation," Applied

Physics Letters 91 (13), 131906 (2007).

8 A. Hu, M. Rybachuk, I. Alkhesho, Q. B. Lu, and W. Duley, "Nanostructure

and sp/sp2 clustering in tetrahedral amorphous carbon thin films grown by

femtosecond laser deposition," Journal of Laser Applications 20 (1), 37-42 (2008).

7 A. Hu, S. Griesing, M. Rybachuk, Q.-B. Lu, and W. W. Duley,

"Nanobuckling and x-ray photoelectron spectra of carbyne-rich tetrahedral carbon

films deposited by femtosecond laser ablation at cryogenic temperatures," Journal of

Applied Physics 102 (7), 074311-074316 (2007).

6 A. Hu, Q.-B. Lu, W. W. Duley, and M. Rybachuk, "Spectroscopic

characterization of carbon chains in nanostructured tetrahedral carbon films

synthesized by femtosecond pulsed laser deposition," The Journal of Chemical

Physics 126 (15), 154705 (2007).

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5 M. Rybachuk and J. M. Bell, "The effect of sp2 fraction and bonding disorder

on micro-mechanical and electronic properties of a-C:H films," Thin Solid Films 515

(20-21), 7855-7860 (2007).

4 M. Rybachuk and J.M. Bell, "Synthesis of diamond-like carbon films using a

bi-modal sputter deposition with Xe ions," Surface Review and Letters 14 (4), 735-

738 (2007).

3 M. Rybachuk and J. M. Bell, "The observation of sp2 fraction disorder using

dual wavelength Raman spectroscopy in a-C:H films fabricated using an open

inductively coupled plasma reactor," Diamond and Related Materials 15 (4-8), 977-

981 (2006).

Refereed conference contributions

2 M. Rybachuk, J.M. Bell “Growth of ta-C films using low energy ion beam

sputter - bombardment deposition with Ar ions”, Journal of Physics: Conference

Series 100, 082009 (2008)

1 M. Rybachuk and J.M. Bell, "The morphology of hydrogenated diamond-like

films and the effect of the sp2 fraction disorder on electronic and micro-mechanical

properties," Proceedings of SPIE, Microelectronics: Design, Technology, and

Packaging II; Alex J. Hariz; Ed. 6035 (603503) (2006).

Other conference contributions

9' M. Rybachuk, J.M. Bell "Growth of ta-C and a-C:N thin films using a bi-

modal sputter deposition with Xe and N ions", Proceedings of 17th International

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Vacuum Congress / 13th International Conference on Surface Science, Stockholm,

Sweden 7/2007.

8' A. Hu, M. Rybachuk, Q. - B. Lu, W. W. Duley “Femtosecond pulsed laser

deposition and optical properties of diamond-like amorphous carbon films embedded

with sp-bonded carbon chains”, Proceedings of DIAMOND – 18th Int. Conference on

Diamond and Related Materials, Berlin, Germany, 9/2007

7' M. Rybachuk, J.M. Bell "Growth of ta-C films using low energy, reactive ion

beam sputter deposition with Ar and Xe ions", Proceedings of DIAMOND – 18th

International Conference on Diamond and Related Materials, Berlin, Germany,

9/2007

6' M. Rybachuk, J.M. Bell “Fabrication of DLC films using a novel bi-modal

ion beam bombardment deposition”, Proceedings of Asia-Pacific Conference on

Surface Science, Hong Kong, 12/2006.

5' M. Rybachuk M., J.M Bell “Phenomenological study of the sp2 fraction

arrangement and sp2:sp3 evolution in of a-C:H films”, Proceedings of DIAMOND –

16th International Conference on Diamond and Related Materials, Toulouse, France

9/2005

4' M. Rybachuk M., J.M Bell “The effect of the sp2 fraction arrangement on

mechanical properties of a-C:H films” and “Disorder in hydrogenated diamond-like

carbon films”, EUROMAT European Congress on Advanced Materials and

Processes, Prague, Czech Republic, 9/2005.

3' M. Rybachuk M., J.M Bell “Systematic analysis of a-C:H films deposited

using open plasma generator”, Proceedings of the 3rd International Conference on

Hot-Wire CVD, Utrecht, Netherlands 7/2004.

2' M. Rybachuk M., J.M Bell “Multimode Raman spectral analysis of a –C:H

films”, Proceedings of the 19th international Conference of Raman Spectroscopy,

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Surface Paradise (Australia), P. M. Fredericks, R. Frost and L. Rintoul (Eds.),

CSIRO Publishing, 8/2004.

1' M. Rybachuk M., J.M Bell “Fabrication of hard a-C:H coatings in an open

plasma”, Proceedings of the 28th Condensed Matter and Material Meeting, Wagga

Wagga, Australia, 2/2004.

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_____________________________________________

Abstract

This thesis presents theoretical and experimental study of properties of

amorphous diamond-like carbon (DLC) coatings synthesised using discharge plasma

methods. There were two objectives in this study.

The first objective was to investigate the formation mechanism of hydrogenated

DLC films (a-C:H) in an open hydrocarbon plasma source. The inductively coupled

plasma (ICP) reactor was used to synthesise the films and the formation of sp2 and

sp3 hybridised phases and the combination of these phases in the ICP plasma

environment was studied. It was found that for a-C:H films with narrow distribution

of the sp3 content the mechanical properties are determined by the degree of disorder

of the sp2 fraction. The relationship between the sp3 content in fabricated films and

hardness and Young's modulus was established. Raman and multi-wavelength (Vis –

UV range) Raman spectroscopy was primarily used together with other suitable

analytical methods to examine a-C:H films and it was found that films fabricated at

higher ion energies displayed higher degree of clustering and bonding disorder than

films produced at lower ion energies. All as fabricated a-C:H films were also found

contain basic π-conjugated polymer inclusions as of trans-polyacetylene. The Raman

results also reveal that the magnitude of Rayleigh scattered light is related to the

relative density of the films, a feature that can be useful for monitoring film growth

in-situ. The use of X-ray photoelectron spectroscopy (XPS) as a suitable method for

measuring the sp3 content of the bulk DLC was also established.

The second objective was to develop a fabrication technique that would allow

fabrication of DLC films using graphite target sputtering with a single focused ion

beam source and producing films with medium-high sp3 content. This research was

motivated by the industrial partner of the project Laserdyne Pty Ltd that required a

simple DLC deposition apparatus to be integrated into a standard, stand alone,

optical thin film deposition chamber. Such technique was developed on the basis of a

conventional ion beam target sputtering. In our experiments hydrogen-free DLC

films with medium sp3 content were produced using a single, Kaufmann type ion

source operated at low energies. The fabrication technique, denoted a reactive ion

beam sputter deposition (RIBSD), was based on sputtering a graphite target at low

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incident angles and positioning the substrate at the grazing angles to the incoming

ions, thus the incident ions (Ar and Xe ions were used) were simultaneously

bombarding the target and the growing film. The effect of angle of incidence of an

ion beam to the target and to the substrate in creating the sp3 content in DLC was

investigated. It was found that the infringement bombardment of the substrate was

not favourable for DLC growth as it essentially provided for a secondary re-

sputtering process. Quality DLC films with approximately 40 % of the sp3 content

were fabricated at the optimal angle of the ion flux to the target of 30º and to the

substrate of 0º (parallel to the ion bema axis). The increased ion energy contributed

to structural changes in DLC from predominantly sp2 graphitic like bonding to

tetrahedral sp3 bonding arrangement.

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Contents ____________________________________________________________________

1. Introduction

1.1 Thesis outline and Scope

1.2 Background on DLC

1.2.1 Open questions

1.2.1.1. Elucidate the formation of hard a-C:H and whether

sp3 bonded atoms control the properties of the films

1.2.1.2. Develop a new DLC fabrication technique

2. Overview of DLC fabrication techniques and DLC formation

mechanism

2.1 Overview of DLC fabrication techniques

2.1.1 Ion beam deposition

2.1.2 Mass selected ion beam

2.1.3 Sputtering

2.1.4 Cathodic arc

2.1.5 Laser ablation

2.1.6 Plasma enhanced CVD

2.1.7 Summary of DLC fabrication

2.2 Deposition mechanism of hydrogen free DLC

2.2.1 Specifics of hydrogenated DLC growth

2.2.2 Types of hydrogenated DLCs

3. Experimental methods used to fabricate hydrogenated and hydrogen

free DLC

3.1 The ICP system used for fabrication of a-C:H films

3.1.1 a-C:H experimental arrangements

3.2 Growth of hydrogen free DLC using the RIBSD system

3.2.1 Monte Carlo simulations of Ar and Xe ions interactions with

a target

3.2.2 The RIBSD experimental arrangements

Page

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4. Analytical methods used to study the fabricated DLC films

4.1 Approach to DLC structure – properties characterisation

4.2 Nanoindentation measurements

4.2.1 Instrumental settings UMIS

4.3 X-ray photoelectron spectroscopy and obtaining sp2/sp3 ratio from

the core level C1s peak

4.3.1 Instrumental settings XPS

4.4 Multi-wavelength Raman spectroscopy

4.4.1 Lineshape dilemma

4.4.2 Rayleigh scattering measurements

4.4.3 Instrumental settings Raman UV- Vis

4.5 Fourier transform infrared spectroscopy

4.5.1 Instrumental settings IR

4.6 Band gap (Tauc gap) and surface conduction gap

4.6.1 Instrumental settings STS

4.7 Scanning electron microscopy

5. Characterisation of fabricated a-C:H films

5.1 Nanoindentation results and discussion for a-C:H samples

5.2 X-ray C1s results discussion for a-C:H

5.3 MW Raman spectroscopy results and discussion for a-C:H

5.3.1 Rayleigh scattering results and discussion

5.4 IR spectroscopy results and discussion for a-C:H

5.5 Band gap measurements for a-C:H. Results and discussion of Tauc

gap vs. surface conduction gap

5.6 SEM images for a-C:H

5.7 Discussion of obtained results for a-C:H

6. Characterisation of hydrogen free DLC fabricated using the RIBSD

6.1 UV Raman analysis for hydrogen free DLC

6.2 X-ray C1s results for hydrogen free DLC

6.3 Discussion of the RIBSD technique for DLC fabrication

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7. Summary of the results and contributions to the existing field of

knowledge

7.1 Main findings of the work on hydrogenated DLC films

7.2 Main findings of the work on development and investigation of a

new deposition technique (the RIBSD)

7.3 Future outlook

References

Appendix 1

Nanoindentation diagrams for the ICP fabricated a-C:H films

Appendix 2

XPS C1s data for the ICP fabricated a-C:H films

Appendix 3

Multi-wavelength Raman spectroscopy data for the ICP fabricated a-C:H

films

Appendix 4

IR spectroscopy data for the ICP fabricated a-C:H films

Appendix 5

UV Raman spectroscopy data for the RIBSD fabricated films

Appendix 6

XPS C1s data for the RIBSD fabricated films

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List of Figures

Chapter 1

Fig.1. Carbon hybridised states sp3, sp2 and sp. From Ref. 7

Fig. 2. Comparison of calculated sp3 fraction of DLC with C ion bombardment

energy. From Ref. 17.

Fig. 3. Ternary phase diagram of composition in C-H alloys. From Ref. 29. The

dotted circle shows the region of the phase diagram where films in this work are

located.

Chapter 2

Fig. 4. Ion beam deposition. From Ref. 1

Fig. 5. Ar plasma sputtering of a graphite target. From Ref. 1.

Fig. 6. Ion assisted sputtering

Fig. 7. Cathodic vacuum arc (CA). From Ref. 1

Fig. 8. Pulsed laser deposition (PLD) system. From Ref.139

Fig. 9. Plasma beam source. From Ref. 1,41.

Fig. 10. Ion ranges and yields of ion processes in carbon. From Ref.1,19

xviii

Fig. 11. Schematics of densification mechanism in a carbonaceous solid by

subplantation. From Ref.1.

Fig.12. Subplantation schematics: direct penetration, penetration by knock-on of a

surface atom and relaxation of a region with higher density. From Ref.1.

Fig. 13. Subplantation schematics by a molecular ion. From Ref.1.

Fig. 14. Comparison of the sp3 fraction of ta-C:H to that calculated by the

subplantation model. From Ref. 1,41,75.

Fig. 15. Schematic diagram of the subplantation process showing a transition of

energy levels. From Ref. 25.

Fig. 16. Berman-Simon phase diagram for carbon. From Ref. 1,27.

Fig. 17. Process diagram of subplantation, when specific interstitial configurations

are included. From Ref. 25.

Fig. 18. Variation of sp3 content with deposition temperature at 100 eV. From Ref. 25.

Fig. 19. Sp3 fraction vs. ion energy at various deposition temperatures. From Ref. 25.

Fig. 20. Growth mechanism of hydrogenated DLCs. From Ref. 1

Chapter 3

Fig. 21. Schematic diagram of open plasma reactor.

Fig. 22. Open plasma source reactor (courtesy of Laserdyne Pty Ltd).

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Fig. 23. Ion beam sputter deposition IBSD (single beam).

Fig. 24. Schematics of RIBSD, where a single ion beam is used for target sputtering

and concurrent substrate bombardment.

Fig. 25. Schematics of relative target and substrate positions to the incident ion beam

flux.

Fig. 26. Percentage of BS ions as a function of the ion energy and the incidence

angle. Ar bombardment of HOPG target. Inset shows the angle definition used for

clarity.

Fig. 27. Percentage of BS ions as a function of the ion energy and the incidence

angle. Xe bombardment of HOPG target.

Fig. 28. The relationship between the energy of the sputtered C atoms, the incoming

ion bombardment energy and the angle of ion incidence. Ar bombardment of HOPG

target.

Fig. 29. The relationship between the energy of the sputtered C atoms, the incoming

ion bombardment energy and the angle of ion incidence. Xe bombardment of HOPG

target.

Fig. 30. Number of C atoms ejected per single Ar ion at a given incident angle.

Fig. 31. Number of C atoms ejected per single Xe ion at a given incident angle.

Fig. 32. The RIBSD experimental set up (courtesy of Laserdyne Pty Ltd). Kaufmann

ion source is shown on the left and the target/substrate holder is on the right.

Fig. 33. The RIBSD in operation. Argon plasma discharge is visible between the

body of the ion gun and the target/substrate holder.

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Chapter 4

Fig. 34. Schematic VDOS of a carbon showing σ and π states. From Ref. 1,198.

Fig. 35. Multiple-point unload method uses the slope of the tangent to the initial

unloading to determine hp. Single-point unloading is faster, and hence it is less

affected by the thermal drift, but relies on a single data point in the unloading portion

of the test cycle (From Oliver and Pharr220).

Fig. 36. The C1s diamond spectrum. From Ref. 235.

Fig. 37. Measured C 1s photoelectron spectra of a a-C film. The Shirley background

and the sp2 and sp3 components resulting from the fit are also shown 241.

Fig. 38. Raman spectra of carbonaceous materials. From Ref.1.

Fig. 39. MW Raman spectra of (a) ta-C, (b) ta-C:H, (c) sputtered a-C and (d)

polymeric a-C:H. The peaks’ trends are indicated. From Ref. 260,262.

Fig. 40 (A, B1, B2). A: The dispersion of the G peak vs. excitation

wavelength/energy for a series of template samples. B1 and B2: I(T)/I(G) and T peak

positions vs. sp3 fraction for non-hydrogenated carbon films. From Ref. 260,262.

Fig. 41. IR spectrum of an a-C:H film. From Ref. 291,292.

Fig. 42. Calculated variation of band gap with sp2 fraction. From Ref. 301.

Fig. 43. Variation of Tauc gap with sp2 fraction 311.

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Chapter 5

Fig. 44. Hardness as a function of penetration depth for samples deposited at -400 V

and -250 V negative bias. The Si <100> substrate hardness (13.2 GPa) and the mean

surface hardness Hn, are indicated by horizontal lines.

Fig. 45. Young’s modulus as a function of penetration depth for samples deposited at

-400 V and -250 V negative bias. The Si <100> substrate E modulus (145 GPa) and

the mean surface Young’s modulus Eio, are shown by horizontal lines.

Fig. 46. Load propagation dP/dh vs. penetration depth h, for a-C:H samples

deposited at different bias.

Fig. 47. The C1s spectra of a-C:H films fabricated under varying substrate bias.

Fig. 48. The fitted XPS C1s spectrum of an a-C:H film fabricated under – 400 V.

Fig. 49. MW Raman spectrum of an a-C:H film fabricated under -400 V bias.

Fig. 50. Relationship between the height of 532 nm scattered Rayleigh line for a-C:H

samples fabricated under different bias.

Fig. 51. IR absorption spectrum of a-C:H sample deposited at -250 V.

Fig. 52. The deconvoluted IR absorption spectra in 2700 – 3200 cm-1 region for a-

C:H samples deposited at -250 and -400 V

Fig. 53. Comparative IR absorption spectra in 1050 – 1700 cm-1 region for a-C:H

samples fabricated at -250 V and -400 V substrate bias.

Fig. 54. Extrapolation of N-IR absorption spectra for a-C:H samples.

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Fig. 55. The variation of the Tauc, Eg and surface conduction gap, ESC with bias

voltage for examined a-C:H films.

Fig. 56. A) Schematics VDOS of DLC and B) Perturbation of π states is shown.

Fig. 57. The surface of an a-C:H film fabricated under -250 V bias.

Fig. 58. The surface of an a-C:H film fabricated under -300 V bias.

Fig. 59. The surface of an a-C:H film fabricated under -350 V bias.

Fig. 60. The surface of an a-C:H film fabricated under -400 V bias.

Fig. 61. Lateral image of an a-C:H film on Si substrate.

Fig. 62. UV Raman spectra of a-C films fabricated at sputtering angles of αt 30° and

αs 10° and sputtering ion energy of 1.2 keV for Ar and Xe ions.

Fig. 63. Fitted UV Raman spectra of an DLC film fabricated using 1.0 keV Xe ions.

The αt, : αs was 30° : 0°.

Fig. 64. The relationship between the ion bombardment energy and I(D)/I(G) ratio

for the RIBSD fabricated films. The legend shows the angles of target and substrate

as αt_ αs

Fig. 65. The relationship between the ion bombardment energy and I(T)/I(G) ratio for

the RIBSD fabricated films.

Fig. 66 A: sp2 rich a-C film fabricated using Ar bombardment at 1.0 keV and αt, : αs

of 45° : 10°, B a DLC film fabricated using Xe ions at 1.0 keV and αt, : αs of 30° : 0°.

Fig. 67 The relationship between the sp3 content (XPS C1s) in the RIBSD films and

the T peak intensity (325 nm UV Raman).

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Chapter 6

Fig. 59. UV Raman spectra of a-C films fabricated at sputtering angles of αt 30° and

αs 10° and sputtering ion energy of 1.2 keV for Ar and Xe ions.

Fig. 60. Deconvoluted UV Raman spectra of a DLC film fabricated using Xe ions

with energy of 1.0 keV. HOPG target and the substrate were positioned at 30° and 0°

(parallel) to the ion beam axis respectively.

Fig. 61. Ion energy for Ar and Xe ions as a function of I(D)/I(G) ratio for fabricated

a-C and ta-C films; from 325 nm UV Raman spectra.

Fig. 62. Relationship between Ar and Xe ion beam sputtering energy and I(T)/I(G)

ratio; from 325 nm UV Raman spectra.

Fig. 63. Ta-C film fabricated using sputter-bombardment with 1.0 keV energetic Xe

ions; the target was positioned at 30° and the substrate was set parallel to the centre

of the ion beam axis.

Fig. 64. Sp2 rich a-C film fabricated using sputter-bombardment with 1.0 keV

energetic Ar ions; the target was positioned at 45° and the substrate was set at 10° to

the centre of the ion beam axis.

Fig. 65. Cross section of ta-C film fabricated using Ar ions with energy of 1.2 keV.

The angle of HOPG target to the ion beam axis was 30° and the substrate was

positioned parallel to the incoming Ar ion beam. Lighter coloured ta-C film is darker

coloured Si substrate. Flake-like appearance of the Si substrate is due to crack

propagation through <111> lattice.

Fig. 66. Frontal surface area of ta-C film fabricated using 1.2 keV Ar ions at 30 : 0

sputter-bombardment arrangement. The resolution is 2 µ.

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Fig. 67. Same surface area of ta-C film as in Fig. 65 at higher magnification. The

resolution is 300 nm.

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List of Tables

Chapter 2

Table 1

Comparison of properties of carbonaceous materials

Chapter 3

Table 2

The ion types used during the RIBSD experiments, their respective energies and the

angles αt and αs (expressed as αt : αs (in °)).

Table 3

Operation variables during RIBSD experiments.

Chapter 4

Table 4

Comparison of characterisation methods available for DLC analysis, their advantages

and availability.

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Chapter 5

Table 5

Nanoindentation results for a-C:H samples produced under various substrate bias.

Table 6

The detailed XPS C1s results for samples produced under various substrate bias. The

fitting uncertainties for hybridised states are ± 0.02 eV, for the sp3/sp2 ± 0.018 and

for ∆BE is ± 0.03 eV.

Table 7

Assignments of a-C:H IR vibrational frequencies in the 2700 – 3200 cm-1 region; C-

H stretching vibrations 344-346.

Table 8

Calculated relative peak areas, A as a % of total peak area for C-H stretching groups

in 2700 – 3200 cm-1 region.

Table 9

The N-IR vibrational frequencies in 1050 – 1700 cm-1 region for a-C:H 343.

Chapter 6

Table 10

Films fabricated using the RIBSD at varying Ar and Xe ion beam energies and

target/substrate sputtering geometry. “--" shows the experiments were no performed.

Table 11

The sp3 content, ± 1.5 % of a-C and DLC films fabricated using the RIBSD method

with Ar and Xe ions.

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Definition of Symbols

∆BE binding energy difference

A area of an indenter

B constant

BEsp2 binding energy for sp2 phase

BEsp3 binding energy for sp3 phase

C Raman cross section area

Ce intercept factor (UMIS standard)

Cf deflection of the load frame

Cs constant accounting for a non-ideal shape of the indenter

dh change in penetration depth

dP change in loading force

E Young's modulus (elastic modulus) of elasticity

E* combined modulus of elasticity

E04 optical absorption band at α = 10-4 cm-1

Eb surface binding energy

Ed displacement threshold energy

Eg band gap

Ei elastic modulus of an indenter (Section 4.2)

Ei ion flux energy (Section 2.2)

Eo diffusion activation energy

Ep penetration threshold energy

Es constant (spread) parameter

Es elastic modulus of a sample

ESC surface conduction gap

F beam of ion flux

H hardness

h penetration depth

hυ photon energy

I current

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I0 peak intensity

Ia axial current

Im magnetic mirror current

Is substrate current

k wavevector

kT isothermal compressibility coefficient

M Pearson width parameter

M number of rings in the cluster (Section 2.2 only)

n fraction of interstitials below the surface

n refractive index

P loading force

PAr pressure for Ar gas

PCH4 pressure for CH4 gas

Q amplitude of a photon (Section only)

Q skewness coefficient for Breit – Wigner – Fano function

q wavevector

r0 is the equilibrium spacing between atoms

T temperature

T0 thermal stability for ta-C (temperature for sp3 phase relaxation)

tanα slope of the unloading curve

Td thermal drift

V potential

Va axial potential

Vm magnetic mirror potential

Vs substrate bias potential

α optical absorption coefficient (Section 4.6)

αs incidence angle of the ion beam to the substrate

αt incidence angle of the ion beam to the target

β relaxation rate constant

γ π – π* interaction gap

∆ ρ density increase

ε optical dielectric constant

ν0 phonon attempt frequency

νi Poisson ratio of an indenter

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νs Poisson ratio of a sample

νsp2 rate for an interstitial transition into sp2 phase

νsp3 rate for an interstitial transition into sp3 phase

ρ density

τ exposure time

φ fraction of an ion beam flux

χ light polarisability

ω photon frequency

ω0 peak position

Г full width at half maximum (FWHM)

Гg FWHM for Gaussian function

Гl FWHM for Lorentzian function

1

Chapter 1

Introduction

This Chapter presents a digest of the work performed during the course of the

project. It starts with the ‘Thesis Outline and Scope’ where content of the each

individual Chapter is presented. Here, a reader is introduced to the world of DLC and

different Sections of this Chapter highlight the research performed as a part of the

project and present a short summary of the obtained results.

1.1 Thesis outline and Scope

This thesis aims to:

1. Discover the mechanism of formation of hard a-C:H coatings and elucidate of

the role of sp3 phase in controlling the mechanical properties of the films.

2. Develop a new deposition method for growth of DLC films using ion beam

sputtering of a graphite target.

Aim № 1 is focused on understanding of the formation mechanism of a-C:H

coatings in hydrocarbon plasmas since these are most widely used in DLC synthesis1,

and understand the structure-property relationship of fabricated DLC coatings. Hard

a-C:H films were selected for the investigation since this class of DLC films have not

been a subject of intensive studies2.

Aim № 2 is to develop a new DLC growth method that does not require

extensive resources to be installed or set up, that is simple to operate, that is robust

and the use of which allows fabrication of DLC films with medium to high sp3

content at reasonable deposition rates. The ease and simplicity of integration of a

2

new DLC deposition method into a conventional optical thin film deposition

chamber is a major interest of the industrial research partner, Laserdyne Pty Ltd 3.

The original DLC growth technique was developed by modification of a

conventional target sputtering method by an ion beam, i.e. ion beam sputtering

deposition or IBSD. This classical abbreviation was modified to RIBSD and, is used

throughout the text to describe our DLC fabrication method. The RIBSD technique is

different from IBSD owing to the use of a single low energy ion source that is

bombarding a target and also simultaneously bombards a substrate thus implying a

reactive condition. The term RIBSD was proposed by Bell et al. 4 and an additional R

denotes 'reactive'. In Section 3.2 we examine how the energies of incoming ions and

the target to substrate geometry affects the formation of the sp3 boding in DLC films.

Two ion species were studied in this part of the project: Ar and Xe.

A review of the relevant literature is presented into Chapter 2 where the theory

of sp3 formation for both hydrogen free and hydrogenated DLC films is discussed. In

Section 2.2 most common DLC deposition techniques and methods are presented.

The experimental methods used in this project to fabricate a-C:H and hydrogen

free DLC films are summarised in Chapter 3. The apparatus, operation and specific

experimental variables are described in detail for the two deposition methods used.

Since specific fabrication parameters for a-C:H in an open plasma system is already

in the public domain5, we only underline several key points of a-C:H fabrication

using this particular system. Monte Carlo simulations using SRIM 6 were used to

model the sputtering process that leads to DLC formation in RIBSD. The SRIM

calculations were performed for Ar and Xe ions bombarding a highly ordered

pyrolytic graphite (HOPG) under variable incident bombardment angles. The results

of these simulations were used to optimise the geometry of the sputtering process and

to predict the likelihood of fabrication of DLC films.

Chapter 4 details analytical methods employed to characterise fabricated a-C:H

and hydrogen free RIBSD produced films. These are: nanoindentation

measurements, X-ray photoelectron spectroscopy (XPS), Raman and the resonant

Raman spectroscopy (RRS), infra-red (IR) and scanning tunnelling (STS)

spectroscopy and scanning electron microscopy (SEM).

3

Chapter 5 presents the results of a-C:H analysis using the analytical techniques

discussed in Chapter 4. Section 5.7 summarises the work and outlines the

contributions to the existing field of knowledge made through this work on a-C:H.

Chapter 6 presents the results of DLC films fabricated using RIBSD. The

deposition variables in RIBSD are discussed and key factors for optimisation of this

technique for fabrication of quality DLC are also proposed.

Chapter 7 summarises the entire work performed on fabrication and analysis of

a-C:H and hydrogen free DLC films and suggests research directions that may be

taken in the future on the topics explored in this thesis.

1.2 Background on DLC

DLC is characterised as a metastable form of amorphous carbon containing a

significant fraction of sp3 bonds1. For thousands of years two pure crystalline forms

of carbon were known and used by mankind: graphite and natural diamond. Other

less illustrious but, nevertheless viable carbonaceous materials are also known, these

are soots, coals and different tars. Carbon forms so many crystalline and disordered

structures because it is able to exist in three hybridised forms, sp3, sp2 and sp, Fig. 1.

Fig.1. Carbon hybridised states sp3, sp2 and sp. From Ref. 7

4

In the sp3 configuration, as in natural diamond, four valence electrons of carbon atom

are each assigned to a tetrahedrally aligned sp3 orbital, making a strong σ bond with

adjacent atoms1. In the sp2 configuration, as in graphite, three of four valence

electrons are positioned along a trigonally directed sp2 orbital making a strong σ

bond in a plane, while the fourth electron lies in a π orbital normal to the threefold σ

plane. This π orbital forms a weaker π bond with a π orbital on one or more adjacent

atoms. In the sp configuration, two of four valence electrons enter σ orbitals forming

a σ bond aligned along the x axis, and the other two enter π orbitals along the y and z

axes1. The first DLC film was produced more than 30 years ago by Aisenberg and

Chabot 8. Following their work on epitaxial deposition of Si, that is production of Si

ions by means of sputtering of Si solid state electrodes in Ar plasma, they reversed

the process and sputtered polar carbon electrodes, and the material they produced

was unlike a conventional C everyone knew. The fabricated film was transparent,

hard and insulating. Initially the term ‘DLC’ was exclusively used for hard and

insulating hydrogen free films9, however at present due to the field amorphous

carbon becoming highly commercialised, the ‘DLC’ term is also used for

hydrogenated and even doped (by N, B, P etc.) carbon alloys. DLC can have high

mechanical hardness and Young’s modulus and chemical inertness similar to natural

diamond, optical transparency, and it is a wide band gap semiconductor2,10-16.

Contrary to the diamond crystal itself that exists in pure crystalline form, DLC

properties are achieved in an isotropic disordered thin film which could be from a

few nanometres to tens of microns thick and most importantly, the structure of DLC

is amorphous and does not display any grain boundaries. Fabrication of DLC is much

cheaper than fabrication of diamond itself. The deposition process which promotes

sp3 bonding in DLC is a purely physical ion bombardment process9,13,17,18. DLC

materials with the highest sp3 fraction are formed using carbon ions with ion energy

around 100 eV, as illustrated in Fig. 2. The formation mechanism is unique for DLC

and the most widely accepted theory belongs to Robertson who coined the term

"subplantation" to describe the mechanism of formation18-21. The process describes

as an incoming C ion being ‘sub-planted’ into the bulk of the film and the DLC film

is growing from the inside out (rather from the surface). The increase of sp3 fraction

in the growing film is due to a metastable increase in density that relates to the

energy of bombarding ions (see Fig. 2).

5

Fig. 2. Comparison of calculated sp3 fraction of DLC with C ion bombardment

energy. From Ref. 17.

The early DLC formation model was taken from nature where compressive stress

and high temperatures are vital to stabilise the sp3 fraction in natural diamond.

Various numerical and analytical simulations performed recently confirmed the basic

idea of the subplantation process22-24. However, only in late 2005 were major

outstanding issues regarding subplantation mechanism resolved by Robertson25.

The sp3 bonding is what gives DLC so many of the beneficial properties of

diamond derived from its strong, directional σ bonds 26. Diamond has a wide 5.5 eV

band gap, the largest bulk modulus of any solid and the highest atom density, the

largest thermal conductivity at room temperature, smallest thermal expansion

coefficient and the largest electron and hole velocities of any semiconductor26.

Graphite, in its pure form with its threefold planar sp2 arrangement has strong intra

layer σ bonding and a weak Van der Waals bonding along π bonds between its

layers. A single graphite plane is a zero band gap semiconductor and in three

dimensions it is an anisotropic metal27,28.

The group of materials known as DLC includes not only amorphous carbons

but also hydrogenated alloys, a-C:H, that can be soft and hard, and tetrahedral a-C:H

denoted as ta-C:H. Jacob and Moller29 proposed a ternary diagram to summarise the

various forms of C-H alloys, as shown in Fig. 3.

6

Fig. 3. Ternary phase diagram of composition in C-H alloys. From Ref. 29. The

dotted circle shows the region of the phase diagram where films in this work are

located.

On this diagram the amorphous carbons with disordered graphitic ordering like soot,

chars, glassy carbon and evaporated a-C lie in the lower left hand corner.

Polyethylene -(CH2)x and polyacetylene -(CH)x define the limits of a triangle in the

right hand corner beyond which interconnecting C-C networks cannot form. This is

the limit where only molecules are formed. Films with tetrahedral sp3 bonding have

been denoted by McKenzie13 as ta-C in order to distinguish them from

predominantly sp2 rich amorphous a-C films. Some deposition methods such as

plasma enhanced chemical vapour deposition method (PECVD) are able to reach the

interior of this triangle 2 fabricating a-C:H and ta-C:H films with high sp3 content.

Fig. 3 shows that the sp3 content in a-C:H films is not high, however the hydrogen

content could be very large. Hydrogen play a vital role in the formation of DLC films

and films with the same sp3 content but with different hydrogen content can have

different properties30-36. The ability to tailor DLC properties is advantageous for

many applications and Table 1 summarises various forms of DLC as compared to

natural diamond, graphite and other types of carbonaceous materials1,2,7,13,17,26-28,37-43.

The applications of DLC materials stem from their unique set of properties that

include high hardness and Young’s modulus13,17,38, excellent wear resistance due to a

low friction coefficient44,45, high thermal stability (with the exception of soft a-C:H),

chemical inertness46,47, and bio-compatibility 48,49.

7

Table 1

Comparison of properties of carbonaceous materials

Material sp3 (%) H- (%) Density,

g cm-3

Eg,

eV

Hardness,

GPa Ref

Diamond 100 0 3.515 5.5 100 26

Graphite 0 0 2.267 0 27,28

C60 0 0 1.6 27,37

Glassy C 0 0 1.3 - 1.6 0.01 3 7

Evaporated C 0 0 1.9 0.4 – 0.7 3 7

Sputtered C 5 0 2.2 0.5 38

ta-C 80 – 88 0 3.1 2.5 80 13,17,38

a-C:H hard 40 30 – 40 1.6 – 2.2 1.1 – 1.7 10 – 20 2,39

a-C:H soft 60 40 – 50 1.2 – 1.6 1.7 – 4.0 < 10 2,40

ta-C:H 70 30 2.4 2.0 – 2.5 < 50 40-42

Polyethylene 100 67 0.92 6.0 0.01 43

DLC films are also exceptionally smooth50,51 with a wide and relatively easy

tuneable band gap52,53, excellent optical transparency in N-IR to mid-UV range54,55,

and good thermal conductance56. Intensive research work performed in the past two

decades eventually paid off when industry incorporated hard a-C:H, ta-C:H and ta-C

films into several applications in the fields of biomedical engineering and optics, but

the most important applications of DLC coatings are in magnetic storage devices57,58.

ta-C:H and ta-C coatings on hard discs eliminate adhesion of the disc to the head

during stop-start events preventing catastrophic failure. This reduces friction on tape-

recording heads and tape transport guides and prevents oxidation of a metal-film

recording tape. a-C:H, ta-C and ta-C:H are often used in wear protective and

antireflective coatings for IR windows59,60. Often a-C:H wear protective coatings are

used in mechanical engineering applications as slide bearings61,62. ta-C and ta-C:H

films are also used to coat precision gauges for the automotive industry. However, a-

C:H coatings have not been successful in application to machining tools as the films

are unstable at temperatures of over 300°C47,63; a-C:H films also display high friction

coefficient in humid environments64,65 and adhesion problems to variety of materials

and especially steels, requiring substantial SiC interstitial layers66,67. As a result these

8

films have not attracted such a widespread use as hydrogen free DLC. The electronic

applications of both a-C:H and ta-C:H have been investigated thoroughly, however

this has only resulted in fabrication of some basic devices at a laboratory level68-70.

At the present, it is unlikely that hydrogenated a-C:H or ta-C:H will re-emerge in

industrial electronics as their electronic properties are controlled by configuration of

the sp2 phase and hydrogenation level, parameters that are difficult to control. The

future electronic applications employing DLC will include nanodiamond71-73 with its

superior properties which are relatively easy to tailor.

1.2.1 Open Questions

1.2.1.1 Elucidate the formation of hard a-C:H and whether sp3

bonded atoms control the properties of the films

The a-C:H films studied in work are situated in the middle to lower part of the

phase diagram shown in Fig. 3 (dotted circle). These films have relatively low sp3

content, often less then 30%. Extensive research performed in the past on a-C:H and

ta-C:H films2,41,42,74,75 has resulted in development of understanding about how

hydrogenated films are formed and several key parameters (ions, ion types, ions

energies and the effect of hydrogen) were identified (see Chapter 2). It is now known

that the increase of the carbon ion energy during the formation process leads a

growing hydrogenated film to pass through several specific morphological transitions

which determine the microstructure and resultant film properties. There are four

recognised1,2 classes of hydrogenated DLC films: polymeric, soft, hard and ta-C:H.

The transition mechanism from polymeric to soft films, then to the hard, and over to

ta-C:H with relaxation back to polymeric as a function of increasing C ion energy,

has been studied extensively in the past76-79. The relationship between the ratio of sp3

hybridised carbon to sp2 hybridised carbon ratio (the sp3/sp2 ratio) and band gap has

been determined14,80 and the early stages of a-C:H nucleation81 have been described.

However, there are only very few reports82,83 where group contributions of each of

9

the sp2 and the sp3 constituents in the narrow range (minor variation of sp3

constituent) have been identified and their influence on a-C:H bulk (mechanical and

opto-electronic) properties fully assigned.

The focus of our investigation will be on elucidation of the role and the effect

of these individual C-H constituent groups in hard a-C:H films. a-C:H films will be

fabricated in a narrow range of fabrication variables (ion energy, ion types) and

microstructure (sp2 and sp3 bonding, bonding arrangement), and the resulting

mechanical and electronic properties of these films will be studied. The aim is to

determine whether the amount of the sp3 bonded atoms in a-C:H films determine the

mechanical properties.

1.2.1.2 Development of a new DLC fabrication technique

Currently the number of available DLC deposition techniques is over two

dozen1,9, however, none of the existing DLC deposition methods can be seamlessly

integrated into a conventional optical thin film deposition chamber (See Section 1.1).

The IBSD84-86 uses a beam of ions to sputter from a graphite target creating a carbon

flux. Often in IBSD process an additional ion beam is used to bombard the growing

film thus delivering the extra ion energy to the densification process. This results in

favourable morphological changes in the growing film (increased density, low

stress)87,88. The RIBSD technique where a single ion beam source in used to bombard

a solid carbon source synthesising an amorphous carbon nitride had been earlier

reported by Bell et al4; an HOPG target was sputtered by using N+ ions and a

substrate was positioned parallel to the central axis of the beam.

The RIBSD technique applied to DLC promises a simple deposition with

minimum of recourses required. The type and quality of fabricated films can be

controlled by selecting ions of certain type, their energy; the target to the ion beam

axis angle will determine the sputtered ion flux such as local plasma density and in

turn establish the film growth rate. Adjusting the angle of the substrate to the ion

beam axis the way where incoming ions bombard or 'plate' the growing film may be

10

found beneficial to the promotion of the sp3 bonding in DLC. The use of the

impinging ion beam on the substrate may provide additional energy to the nucleation

process thus eliminating the need for a secondary ion beam. This work will

investigate how positioning of the target to the ion beam affects the formation of

DLC films (sp3 bonding) and, whether positioning of the substrate at the grazing

angles is beneficial to the promotion of the sp3 bonding. Two ion types will be used

in our RIBSD experiments: Ar and Xe.

11

Chapter 2

1. 2 Overview of DLC fabrication techniques and

DLC formation mechanism

It is important to understand formation of DLC in order to synthesise films

with predetermined properties. In this Chapter several DLC fabrication techniques

that are currently used in fundamental research and in industrial fabricating facilities

are discussed and the formation mechanism is presented the way it was evolving

from the time when the first DLC films were discovered89.

2.1 Overview of DLC fabrication techniques

The first DLCs films were fabricated using ion beam deposition8. Nowadays,

after over three decades of intensive research, numerous deposition methods were

developed which are either producing DLC films suitable for laboratory research or

tailor-make films for large scale industrial production. The general feature of all

fabrication methods is that quality DLC films are formed from a beam of C+ or C-H

ions with the energy of approximately 100 eV. There is a physical deposition process

where an impact of these 100 eV charged ions results in a growing ta-C or ta-C:H

film with predominantly sp3 bonding1,13,90,91. A contrasting approach to a purely

physical deposition is chemical vapour deposition (CVD) of diamond, a-C:H and ta-

C:H films where a chemical process stabilises the sp3 bonding92-95. In this Section

such deposition systems are methods are discussed.

12

2.1.1 Ion beam deposition

In a typical ion beam deposition system, carbon ions are produced by the

plasma sputtering of a graphite cathode in an ion source 1,8,26,96-100 as shown in Fig. 4.

Fig. 4. Ion beam deposition. From Ref. 1

If a Kaufmann type ion source101 is used then hydrocarbon gas is ionised to form a

plasma beam101,102. An ion beam is extracted through a grid from the plasma source

by an externally applied bias voltage. The ions are then accelerated from the grid

towards the substrate forming an ion beam in the medium or high vacuum deposition

pressure. The ion source is operated, however at a finite pressure therefore the beam

always contains a large flux of un-ionised neutral species (C or C-H atoms). This can

reduce the ion/neutral flux ratio to as low as 2-10%1. Conventional Kaufmann type

ion beam sources are most efficient when operated within 0.15-1.0 keV energy

range103.

2.1.2 Mass selected ion beam

High quality research work demands a controlled deposition process where

there are only selected ion species and at predetermined energies. This is achieved by

a mass selected ion beam process (MSIB)9,104-112. In MSIB carbon ions are produced

in an ion source from a graphite target, with a narrow spread of ion energies (less

than 10 eV). Ions are then accelerated to 5-40 keV and passed through a magnetic

13

filter. The filter separates any neutrals from the ions and selects only C+ ions. Due to

Coulombic forces C+ ions are diverged therefore they have to be magnetically

focused. Then they are decelerated to the desired ion energy by electrostatic lenses

and, the resultant, a very uniform C+ beam is focused on a substrate in a high vacuum

producing ta-C films. The advantages of MSIB are obvious due to controllable

deposition process (ion species, energy). However, low deposition rates of 0.0001 Å

s-1 and a high cost and the size of the MSIB apparatus reserve this type of DLC

fabrication to few research laboratories only1.

2.1.3 Sputtering

The most common industrial process for the fabrication of DLC is

sputtering4,113-124. Sputtering in its most widespread form uses either a direct current

(DC) or a radio frequency (RF) sputtering of a graphite target electrode by Ar

plasma, as illustrated in Fig. 51. Since the sputter yield of graphite is low the

enhancement is used in a form of a magnetron sputtering (MS). In MS arrangement

electrostatic magnets are placed behind the sputter target thus causing the electrons

to spiral around and increase their path length. As a result, the degree of plasma

ionisation is also increased.

Fig. 5. Ar plasma sputtering of a graphite target. From Ref. 1.

14

Hydrogenated DLC films could be produced using a reactive sputtering where

plasma is composed of a mix of Ar and H atoms or, for example, a hydrocarbon gas.

Amorphous carbon nitrides (a-C:N) and hydrogenated carbon nitrides (a-C:H:N)

could be produced by using Ar and nitrogen or Ar, N2 and a hydrocarbon gas. A

generic term ‘sputtering’ is also applied to the IBSD (Section 1.2.1.2), and Fig. 6

illustrates an ion beam assisted deposition (IBAD) or ion plating method. Sputtering

is a preferred DLC deposition method for industrial applications because of its

versatility and the ease of scaling up. The sputter deposition conditions can be easily

controlled by the plasma parameters (density, charge type/concentration, ion energy)

and they are independent of the substrate geometry. The main disadvantage of

sputtering lies in a low ion/neutral ratio, the reason why it is practically impossible to

fabricate very high quality DLC films.

Fig. 6. Ion assisted sputtering

However, sputtering methods with a very high fraction of ions have been already

developed 125,126 producing DLC films with a high sp3 fraction but at the expense of a

low growth rate.

15

2.1.4 Cathodic arc

This method utilises the electric discharge between a carbon electrode and an

anode, as illustrated in Fig 7, to initiate a pure carbon plasma thus avoiding

contaminations by Ar or hydrogen as in, for example, sputtering, or ion beam

methods10,127-133.

Fig. 7. Cathodic vacuum arc (CA). From Ref. 1

The source can be operated in a pulsed or DC mode and deliver very high fluxes of

pure carbon (1017 – 1018 atoms cm-2s-1) 9 and high deposition rates of 100 nm min-1.

The ignited plasma mainly consists of carbon ions with a small fraction of neutrals

and microscopic carbon particles. The latter are the reason why DLCs fabricated in

arc discharge are of inferior quality containing graphitic inclusions from tens of

nanometres to few microns in size. Average energy per carbon atom is achieved

about 30 eV9, therefore additional biasing is required to fabricate DLCs with a high

sp3 content. The main advantages of the system are that it enables easy deposition on

insulating substrates and provides uncomplicated doping if needed1.

16

2.1.5 Laser ablation

In this type of DLC deposition laser energy is used to ablate energetic carbon

species from a graphite target 13,134-138, Fig. 8.

Fig. 8. Pulsed laser deposition (PLD) system. From Ref.139

The resultant carbon plume consists of neutral carbon atoms, ions with energy

distribution of up to 70 – 80 eV and graphitic micro-particles. The kinetic energy of

the plume is very similar to MSIB or CA deposition and the DLC films produced are

similar to the films fabricated these techniques. The advantage of PLD is that it is a

versatile method, easily scalable and suitable for deposition of many different

materials. It is a well understood method has been reviewed more than a decade ago

by Voevodin and Donley140.

2.1.6 Plasma enhanced CVD

The most widespread laboratory deposition method is RF plasma enhanced

CVD method (PECVD)92,141-146. The system consists of two electrodes positioned in

different areas. The RF power is usually coupled to the smaller electrode on which

17

the substrate is mounted and the RF power produces plasma between the electrodes.

For DLC deposition, the plasma should be operated at the lowest possible pressure in

order to maximise the ion/radical ratio. However, even at 50 mTorr pressure, the ions

contribute to only about 10% of the film forming flux 1. Therefore, it is necessarily to

use lower pressures but it is not possible for a conventional PECVD as the plasma

would be difficult to initialise. A lower pressure plasma can be created by using a

magnetic field to confine the plasma, and to increase the electron path lengths and

thus increasing the ionisation efficiency. This allows a coupled plasma to operate at

the pressure levels of 5 x 10-4 Torr. At this pressure the ion mean free path exceeds

the sheath thickness and the ion are now distributed with a very narrow energy

margin. This principle became a foundation for the plasma beam source (PBS) 41,

shown in Fig. 9 1.

Fig. 9. Plasma beam source. From Ref. 1,41.

The RF power is applied to a movable electrode so that electrode acquires the

positive self bias. This repels the positive ions through the grid thus forming a

plasma beam which then condenses on the substrate to form a DLC film.

18

2.1.7 Summary of DLC fabrication

This short review of most widely used DCL fabrication techniques aims to illustrate

that all DLC films are formed via the similar mechanism, irrespective of the specific

deposition process. The existing variety of DLC types and their properties relate to

the variation of common set of the deposition parameters for a given deposition

scheme.

These are:

- the distribution of the energetic species (ions, neutrals, hydrogen and doping

species)

- the energy of the species

- the ambient pressure during deposition

- the substrate temperature

- the deposition rate

These parameters are usually pre-determined for a specific DLC to be produced.

2.2 Deposition mechanism of hydrogen free DLC

It has taken more that 20 years to explain the formation mechanism DLC since

first films were produced. Still there are some minor stages of formation phenomena

that remain ambiguous (such as sp3 formation using ultra short laser pulses from the

sp2 phase, the sp to sp3 conversion, etc.) thus warranting the continuation of the

discussion on the DLC formation. The mechanism itself was developed as follows1.

In 1976 Spencer et al.147 proposed various DLC formation mechanisms and one of

his ideas was to suggest that the sp3 phase in DLC arises from the mixture of the

sp3/sp2 sites by a preferential sputtering of the sp2 phase. This idea existed alongside

of the model proposed by Weissmantel et al84,148 who thought that the sp3 phase

arises from a shock wave of the displacement spike of the ion cascade. In early 1990

Lifshitz et al19 noted that Spencer’s process was not likely to work due to sputtering

yield of carbon, and supported his own claim using a comparison drawn theoretically

19

and empirically between the ion energies and yields for various processes of C and H

ions in carbon, illustrated in Fig. 10.

Fig. 10. Ion ranges and yields of ion processes in carbon. From Ref.1,19

The sputtering yield depends mainly on the cohesive energy of an atom, which

is almost the same for the sp2 and sp3 atoms. Therefore, there could be little or no

difference between sputtering yields. Lifshitz et al 19 was the first to notice that the

growth of DLC was sub-surface and the process was denoted “subplantation”. It their

work19 it was proposed that the sp3 phase is formed by preferential displacement of C

atoms of the sp2 phase. However Lifshitz’s concept was flawed as the displacement

threshold of graphite as compared to diamond were found to be very similar149,150.

McKenzie151 and Davis152 together noted that in order to promote formation of the

sp3 phase a compressive stress would be required and this stress could be achieved

by ion beam bombardment. In all these models discussed above preferential

displacement19 on the atomic level was not needed to promote sp3 formation. Only a

subsurface growth in a restricted volume (a derivative of the compressive stress) is

required. Robertson et al18,20,21 suggested that the subplantation process creates a

metastable increase in density causing the local bonding in the sp2/sp3 mixture to

change to sp3. His suggestion was proved to be viable by many workers22-24. As

20

recently as late 2005, Robertson25 again refined the subplantation model that was

proposed earlier by introducing a role of interstitials as independent entities that are

critical to the formation of either sp2 or sp3 phases.

We will consider the formation process on the atomic scale and discuss the

latest proposition25 in more detail. In the energy range of 0.01 – 1 KeV C ions have a

very small range of movement of a few nm1,6. Their energy is lost mainly by elastic

collisions with the target atoms. Assume for the sake of simplicity that all collisions

are binary. At some energy the ions will be able to pass through the surface layer and

we will call that energy the penetration threshold energy, Ep. In order for the ion to

displace an atom from a bonded site, and to create a permanent vacancy-interstitial

pair, the ion has to have a minimum energy and that energy is denoted the

displacement threshold energy, Ed. 1,25

Therefore, the net penetration threshold for free ions could be expressed as

Ep ~ Ed – Eb, (1)

Where Eb, is the surface binding energy. Eb is equal 7.44 eV for C, whereas Ed is

equal 25 eV 1,6.

With the Eq. (1) in mind let us now consider C ions that are incident on an

amorphous carbon surface20,21. An ion at the low energy will not penetrate the

surface and it is likely just to ‘stick’ to the uppermost surface1, and an atom will

remain at its lowest energy, that is an sp2 hybridised state. If the incident ion has a

higher energy than Eb it has high probability to penetrate the surface and enter a

subsurface interstitial site, thus increasing a local density1. The local bonding then

adjusts itself around that atom according to the new density. In amorphous solids,

atomic hybridisations will adjust easily to changes in the local density by becoming

preferentially either more sp2 if the density is low, or sp3, if the density is high1. If

the ion energy increases further there is again a high probability that the ion will

penetrate deeper into the solid. Therefore, at Ed of 25 eV 1,6 for carbon, the Ep will be

about 32 eV from Eq. (1). The Ep value is rather small considering ion energies that

are usually used in DLC deposition. Thus the remaining fraction of the ion energy

will become dissipated in atom displacements153 and the excess of the energy will be

dissipated as photons1.

The whole process consists of three stages1,20:

21

1. a stage of ion collision with a surface, on a time scale of 10-13 s,

2. a stage of thermal dissipation, on a time scale of 10-12 s,

3. a relaxation stage, on a time scale of 10-10 s.

Stages 2 and/or 3 allow the excess density to relax to zero and cause the loss of sp3

phase at high ion energies, see Fig. 2. We need to note that there is a reduction of

density with the increase of ion energy above 150 eV. Let us now consider a beam of

ion flux F, with a fraction of φ, of energetic ions having energy Ei, Fig. 111.

Fig. 11. Schematics of densification mechanism in a carbonaceous solid by

subplantation. From Ref.1.

If a fraction of f of a beam flux φ penetrates the surface, the non energetic

fraction of atoms or ions (1 - f φ) will stick on the outer surface area. Some fraction

of the ions that have already penetrated the surface will relax back to the uppermost

surface. The flux F is proportional to a driving force, which is the fraction of

interstitials below the surface, n. Therefore, the fraction of ions remaining at

interstitial sites and promoting densification will be n = f φ – βn, where β is a

constant1. This gives

βϕ

+=

1fn (2)

Therefore, a fraction n of the ion beam would become ‘sub-planted’ inside the film

and a fraction of (1 – fϕ) is left on the surface as sp2 sites1.

22

The density1 increment of the subplanted fraction is

n1

n−

=ρρΔ (3)

Which gives

βϕϕ

ρρΔ

+−=

f1f (4)

Where ρ is the density of sp2 carbon and Δ ρ is the density increase.

To illustrate an ion penetrating the surface Fig. 12 was used1. It shows that ion

penetration can occur in two ways: directly or indirectly by knock-on.

Fig.12. Subplantation schematics: direct penetration, penetration by knock-on of a

surface atom and relaxation of a region with higher density. From Ref.1.

We note that only knock-on penetration is indicative of ion assisted deposition.

The penetration probability, f could be estimated as a function of ion energy6 and

could be expressed approximately as

⎟⎟⎠

⎞⎜⎜⎝

⎛ −−−=

s

p

EEE

exp1f (5)

23

Where Ep is the penetration threshold, see Eq. (1), and Es is a constant (spread)

parameter1. The first numerical models proposed by Robertson et al20,21 assumed that

relaxation occurs during the thermal spike stage of ~1012 s. That gives a relaxation

rate of β ≈ 0.016 (Ei/Eo)5/3, with Eo as a diffusion activation energy1, and

3/5oi )E/E(016.0f1

f+−

=ϕ

ϕρρΔ (6)

Fig. 3 shows that Eq. (6) gives a good representation of the variation density or sp3

fraction for ta-C with Eo = 3.1 eV. It shows that rising sp3 fraction at low ion energy

is controlled by the penetration probability f, and the decline of sp3 fraction at high

ion energy is controlled by the relaxation process. The diffusion activation energy E0

is derived from the thermal stability, T0 of ta-C exceeding which the relaxation of the

sp3 phase is observed; this temperature is at least 1000 ºC154. Using

E0 = kT0 log(ν0 τ) (7)

Where ν0 is the phonon attempt frequency (~1014 Hz) and τ is the temperature

exposure time1. For τ of, say 1 s, E0 ≈ 3.4 eV. This is a credible value for the self-

diffusion energy of a carbon atom in ta-C, as it is similar to the bond energy of 3.7

eV1.

The penetration/thermal spike model can also account for density dependence

found in hydrogenated DLC (ta-C:H deposited from acetylene)1. The fragmentation

of an incident molecular ion into two C atoms upon the impact with the surface is

shown in Fig. 12. In this figure it is assumed that the ion’s kinetic energy is shared

equally between the constituent C atoms. If we consider a hydrocarbon ion, and a H-

C ion impacting upon a surface as in Fig. 13 then, due to the conservation of

momentum law, hydrogen atoms would retain a minute fraction of the total energy of

the C-H ion. The penetration and densification processed in bulk will occur for each

C atom separately and independently according to Eq. (6)1. A small fraction of

energy should be subtracted from Ei in order for example, to break the C≡C bond in

acetylene. The relaxation step will occur as a single event for the molecular H-C ion

composed of two thermal spikes from the two C atoms.

24

Fig. 13. Subplantation schematics by a molecular ion. From Ref.1.

This could be described as the total energy of the C-H molecular ion.

Therefore, the graph shown in Fig. 14 depicts the energy of a single C atom and the

bulk density of the solid (that is related to the amount of sp3 phase1) as a function of

the incident carbon ion energy. The graph shows a much sharper decline in

densification (sp3 fraction) at higher energies in particular within 100 – 150 eV range

compared to hydrogen free DLC where the incident ion is mono-atomic.

Fig. 14. Comparison of the sp3 fraction of ta-C:H to that calculated by the

subplantation model. From Ref. 1,41,75.

25

Robertson's model, describing subplantation and relaxation back to the initial

state as shown in Fig. 11, could equally be represented in a two-state system1. The

two states corresponding to sp2 and sp3 sites are shown in Fig. 15.

Fig. 15. Schematic diagram of the subplantation process showing a transition of

energy levels. From Ref. 25.

The subplantation process drives C atoms from sp2 to sp3 phase, while

relaxation processes allow them to return from sp3 to sp2 over an energy barrier as

shown in Fig. 15. While there is the evidence that this model is working and it binds

together the theory and the experiment, still the model25 is still deficient as it fails to

account for:

1. the different dependence of sp3 fraction on ion energy in the cathodic arc

(CA) deposition and in the mass selected ion beam (MSIB) deposition. MSIB

gives much slower DLC growth rate compared to CA13,97

2. the transition temperature to sp2 bonding is 400-500K 106,108,155-157 despite the

fact that the temperature in a thermal spike is 106 K

3. variation of the transition temperature for sp3 formation with ion energy22,157

4. variation of the transition temperature with the instantaneous growth

rate158,159

There are many faults with the thermal spike concept above when applied to

DLC deposition. The model25 is only valid for much heavier ions than C ions at high

ion energies, and where the energy loss in the forming solid as a function of distance

(stopping power) is much higher, i.e. at higher energy density. The spike volume

consists of few excited atoms in a site consisting of a network of much less excited

26

atoms, since energy distribution into equal portions is prohibited in carbon1,25. Since

this model was proposed a number of researchers reviewed and re-defined carbon

formation model to account for the inconsistencies noted above. Their suggestions

are outlined below.

Hofsass et al109,153 suggested that the sp3 fraction varies with the number of

hops per atom within a spike. The thermal spike model proposed by Robertson

postulates that there is a driving force of penetration forcing densification and hence

formation of sp3 bonding. Indeed, it becomes ineffective at high ion energies due to

relaxation. Hofsass states that penetration itself has no major role in affecting the

formation of sp3 bonding and there is no driving force but a ‘relaxation’ from an

undefined state into sp3 bonding allowed by the spike.

Hirvonen et al159 noted that DLC tends to possess a range of activation energies

for transport that could be expressed using relaxation behaviour, and Koskinen et

al158 suggested that the dependence of the transition temperature on the DLC growth

rate implies that there must be some overlapping spikes. The incorporation of these

two propositions into the Robertson’s model was not successful. The reason is that

there would be gross overestimation of the size of the site where the relaxation takes

place at transition temperatures. McKenzie et al13,151,160 and Davis152 both proposed

the formation of sp3 bonding is due to presence of compressive stress, and most

importantly the emphasis was on the magnitude of the stress. Their suggestion was

based on the idea that amorphous carbons always exists in ‘quasi-thermodynamic’

equilibrium - the stability of sp2 and sp3 bonding in amorphous carbons follows the

phase diagram of crystalline carbon, shown in Fig. 16. A minimum pressure or

compressive stress above the Berman-Simon line is all what is needed to stabilise sp3

bonding. Again, this idea was not satisfactory as the entire deposition process is

clearly a non-equilibrium process and only some parts of it can be described using

thermodynamic principles. Also the stress model completely forbids the observed

change from the sp3 to sp2 on the premise that these phases remain permanently in a

quenched state. Nonetheless, there is a link between the amount of the sp3 fraction

and stress142,161,162, and molecular dynamics stimulations of networks also do find

that high sp3 networks tend to be under compression163.

27

Fig. 16. Berman-Simon phase diagram for carbon. From Ref. 1,27.

Most recent works by Keriles et al164,165 using molecular dynamics simulation found

the distribution of local stress for each atomic site. Indeed, it was found that sp3 sites

tend to be under compression and sp2 under tension.

Finally, Roberson’s initial model, examined above, for the mechanism of sp3

bond formation was modified by Robertson25 to include a proposed role for

interstitials as independent and explicit species of higher total energy. The interstitial

species are equivalent to an interstitial in a radiation enhanced diffusion process, and

this newest25 model proposes to include this process in the mechanism of the sp3

formation. An interstitial becomes either a sp3 or sp2 site by passing over a small

potential barrier during the energy distribution cascade. sp3 formation occurs by

densification as was previously thought. And the sp3 sites can also anneal or relax

into sp2 state by passing over a large potential barrier.

Fig. 17. Process diagram of subplantation, when specific interstitial configurations

are included. From Ref. 25.

28

In detail Robertson’s most recent proposed model25 suggests that the sub-

planting C ion maintains an interstitial configuration of higher energy during the

cascade process, but not during the thermal spike. The interstitials could be

represented schematically as shown in Fig. 17 by a local minimum at higher energy,

surrounded by two small maxima25. The interstitial must pass over a maximum to

reach either the sp2 or sp3 states 25.

The rate of these processes is given by the following equation for sp2 state

expressed as

νsp2 = ν0 exp( - E2/kT), (8)

and for sp3 state expressed as

νsp3 = ν0 exp( - E3/kT). (9)

The ratio of these rates should depend on the driving force or the local density25. This

can be expressed as

νsp3 / νsp2= exp( - (E3 - E2) kT). (10)

This gives the net transfer rate into sp3 sites25.

We need to point the fact that the relaxation during the thermal spike occurs by

excitation from the sp3 state over the 3.4 eV barrier. For deposition at high

temperatures, the rates described using Eq. (8) and Eq. (9) would definitely compete

with rates driven by the processes of thermal activation. This will cause the rate of an

interstitial turning into the sp2 site to dominate above a transition temperature if E2 ~

kTc. Therefore, E3 should be approximately equal to 0.05 eV. These rates depend

only on the temperature and not the ion flux, F. While all other processes would

remain proportional to flux, F.

Using Eq. (6) and Eq. (9) the final expression for sp3 fraction becomes

29

⎟⎠⎞

⎜⎝⎛+

⎟⎟

⎠

⎞

⎜⎜

⎝

⎛⎟⎟⎠

⎞⎜⎜⎝

⎛+−

=

kTE

expEE

016.0f1F

Ff

30

3/5

0

i νϕ

ϕρρΔ (11)

This expression 25 gives a critical temperature of the correct order, and a transition

that varies with growth flux. Calculated for the most recent model25, Fig. 18 shows

the variation of sp3 content with deposition temperature at 100 eV and Fig. 19 shows

the calculated variation of the sp3 content compared to the ion energy for various

deposition temperatures.

Fig. 18. Variation of sp3 content with deposition temperature at 100 eV. From Ref. 25.

Fig. 19. Sp3 fraction vs. ion energy at various deposition temperatures. From Ref. 25.

30

In these calculations Robertson25 used p = 0.25, Ep = 20 eV and Es = 50 eV in (5),

and E3 = 1.2 eV and v0/F = 1013, or F = 1 atom/unit area/s and v0 = 1013 s-1 in Eq.

(11). These values have been fitted to give the observed Tc. So far this model for the

sp3 formation mechanism has been able to address all existing inconsistencies and

deficiencies that carbon community faced over past two decades.

2.1.1 Specifics of hydrogenated DLC growth

Hydrogenated DLC such as a-C:H, ta-C:H films are deposited from a

hydrocarbon gas source (CH4, C2H2, C2H4 and C6H6) or using sputtering of a carbon

target in an atmosphere filled with hydrogen. Properties of plasma fabricated films

and most importantly the density of fabricated DLC depend strongly on the bias

voltage indicating that the energy of the incident C-H ions plays a critical role in the

deposition. The mechanism of sp3 densification and film growth for hydrogenated

DLC can be described as follows1. As for the hydrogen free DLC films at the onset

we need to assume that an energetic molecular ion that is incident at the film surface

will break up into atomic ions and the energy will be distributed evenly1. Therefore,

each atomic ion will subplant independently with that share of the total energy.

Starting from about 15 – 20 years ago researchers29,166-171 made an attempt to

describe the chemical process involving neutral species and the process of de-

hydrogenation, as well as to describe the physical process of ion subplantation. In

general, their findings29,166-171 can be summarised into a view that there are three

stages of plasma deposition. The first is the reactions in plasma. The second is

reactions between plasma and a surface and, finally the third stage is subsurface

processes in the film that are affected by plasma.

We briefly examine all of them1.

1. The reactions in the plasma.

The plasma reactions are primarily driven by the energetic electrons. However there

are other species formed in secondary reactions that occur in plasma, and example,

31

the species that come to existence due to plasma polymerisation such as polymerised

acetylene and polyacetylene molecules.

2. The plasma-surface interactions.

Plasma species incident on the growing film will consist of ions and neutrals. The

neutrals will be closed shell molecules, single carbon radicals as CH2• and CH4

•,

double-carbon radicals and other unsaturated species such as C2H4• or C2H2

•1. The

plasma will also contain significant amounts of atomic hydrogen, H'. It is known that

neutral species contribute to growth of the film since the mass deposition rate

exceeds the rate due to ions alone. This phenomenon was first noted when the growth

rate decreased with the increasing temperature166. Now it is know that this is due to

etching of the film by atomic hydrogen1,170. The contribution of each of the neutral

species to the growth rate depends on their ability to form a bond and a ‘sticking

coefficient’ (SC)1. Also it is known that a-C:H or ta-C:H surfaces are chemically

passive as it is essentially fully covered by C-H bonds. Double-carbon radicals and

unsaturated species can insert directly into the surface of C-C or C-H bonds and their

SCs approach 11. In contrast, a closed shell neutral like for example CH4 has a very

low SC well below 10-4 thus making a bond creating event negligible. Single carbon

radicals have a moderate effect as they are not able to insert themself directly into the

bond and they only react with the film if there is an existing dangling bond (DB) on

the surface, as shown in Fig. 201.

Fig. 20. Growth mechanism of hydrogenated DLCs. From Ref. 1

32

Single carbon radicals will add to that bond to form C-C bond. However, the

dangling bond could be only created by removal of a hydrogen atom from a C-H

bond1. This can happen in two ways. First, an ion can displace an H from the bond

thus creating a carbon ion with a DB and the H2 molecule. Or as atomic hydrogen H`

abstracting H from the bond also creating a DB and H2 as 1

≡C-H + H` → ≡C` + H2 (12)

Or a DB could be created by another radical like for example CH4 that abstracts H

from the C-H bond and also creates H2 in the process as in

≡C–H + CH4 → ≡C – CH3 + H2 (13)

Neutral hydrocarbon species can only react at the surface level and they cannot

penetrate the film. H` and H ions are different. H` due to being so small can penetrate

about 2 nm into the film1. There, they can again abstract H from C-H bonds and

create subsurface DBs and H2 molecules that are ‘locked in’ within the bulk of the

growing film. Some of the DBs may become re-saturated by incoming H`.

3. Subsurface reactions in the film1.

Ions can penetrate the film and C and C-H ions can cause subplantation. In a-C:H

and ta-C:H ions tend to displace H from C-H bonds as explained above. The

displaced H can then recombine with other H` to form H2 molecule and desorb from

the film or, in a very unlikely event, can become part of an embedded domain entity

as a free gas, see Fig. 201

. This is the main process which causes H content in some a-C:H films to decrease

with increasing bias voltage172-175. Some of the H` does not recombine, but finds DBs

to re-saturate1. Because of their low mass, H` ions interact weakly with C atoms.

Therefore, H+ ions have the longest range and penetrate deepest into the growing

film. They undergo the same reaction as H`, but to a greater depth. Thus, von

Keudell et al176-179 summarised for a-C:H and ta-C:H films to have three

characteristic depths:

33

a) The surface itself. It is controlled by reactions of hydrocarbon and H` and

H species.

b) The upper 2 nm. The chemistry of this layer is controlled by reactions of

H`.

c) The larger depth deeper that 5 nm. This depth depends primarily on ion

energy in which reactions are controlled by H+ ions.

So far the formation mechanism for a-C:H and ta-C:H, and the roles of various

species during the formation, have been confirmed by numerous theoretical and

experimental works, and it is unlikely that the processes described above will be

radically adjusted in the near future.

2.2.2 Types of hydrogenated DLCs

There are four distinct types or classes of hydrogenated DLC films as described

by Koidl et al2, Weiler et al41,42 and by other workers using similar DLC deposition

systems75,145,173,180,181 to the one used in this work, and these film types are:

polymeric, soft, hard and ta-C:H. The formation of a selected type of a-C:H film

whether it is polymeric a-C:H, soft a-C:H, hard a-C:H or ta-C:H is determined by the

energy of C and H ions and neutral species during fabrication process of the film.

Polymeric a-C:H films are formed when the energy of C ions falls well below or well

above 100 eV per single C ion. The result is a mainly sp2 bonded film with either

aromatic or olefinic arrangement and only small amounts of sp3 phase (less than 3%)

present together with a very high hydrogen content (above 70%). sp2 polymeric a-

C:H films are extremely soft, with bulk density between 1.0 – 1.5 gcm-3 and with

virtually no band gap, similar to natural graphite. There are also polymeric a-C:H

films composed of predominantly sp3 bonding with only small amounts of sp2

bonding and a very high hydrogen content of above 60%. They are structurally

similar to polyethylene1. The sp3 polymeric a-C:H films are usually produced from a

low density hydrocarbon plasma with a low degree of radical dissociation. There

were no applications of either these polymeric a-C:H films in the industry due to

34

their inferior properties. Soft a-C:H films are formed as the energy of C ions

increases to low tens of eVs and the process of de-hydrogenation and densification of

the sp3 phase takes place. These soft a-C:H films are diamond-like with sp3 phase

amounting up to 40%2. The sp2 constituent of these films has a higher degree of

bonding disorder with more olefinic sites. Soft a-C:H films display very wide band

gap of up to 4 eV making them useful in opto-electronic applications. However, low

density and poor mechanical properties of soft a-C:H films are the negative

characteristics of this type of the films making their applications limited. Hard a-C:H

films are formed at C ion energies in the at half or below the 100 eV threshold, and

these films are progressively more diamond-like with lesser hydrogen content and

superior mechanical properties than soft a-C:H. Hard a-C:H films display the band

gap up to 2 eV, hardness up to 20 GPa and bulk density of up to 2.2 gcm-3. The sp3

constituent in hard a-C:H films is usually below 40%. Hard a-C:H films are used

widely in industrial applications due to their excellent optical and protective

properties44,182. When the energy of C ions during the deposition process is around

100 eV hard a-C:H films become increasingly tetrahedral, leading to ta-C:H, with sp3

fraction as high as 75 %41,42,75 and hydrogen content as low as 20%. Further ion

energy increase results in relaxation of the sp3 phase as excess energy dissipates as

phonons, resulting in the falling slope of the curve in Fig. 14 above 120 eV. In

addition, the sp3 fraction begins to decrease while H content rises.

35

Chapter 3

3 Experimental methods used to fabricate

hydrogenated and hydrogen free DLC

In this chapter two techniques are presented: the inductively coupled plasma

(ICP) plasma system used to fabricate a-C:H films and the RIBSD technique used for

hydrogen free DLC films synthesis. The design, operation and capabilities of both

systems are described in detail and the experimental particulars are presented.

3.1 The ICP system used for fabrication of a-C:H films

a-C:H films were fabricated using the ICP plasma reactor designed by a team

of researchers from the Australian Defence Science and Technology Organisation

lead by Varga5. Fig. 21 shows the schematics of the system and the original

photograph of the system is presented in Fig. 22.

Fig. 21. Schematic diagram of open plasma reactor.

36

In its simplest form the ICP reactor consists of a vacuum chamber, a cylindrical

water-cooled anode that envelops a helical filament, two gas inlets, and two

Helmholtz type electromagnets with power supplies. The electromagnets provide a

magnetic mirror that confines the plasma within the workspace and one of the

magnets serves as a part of the primary ion source. The helical filament, the anode

and two electromagnets are arranged on a common axis with the filament positioned

within the magnetic mirror. The chamber is evacuated to a base vacuum of

approximately 10-6 Torr range and a common operating vacuum level is

approximately 10-4 Torr. Plasma generation occur inside the cylindrical anode where

electrons are confined by the magnetic field. This is a standard type of an ion source

that is often used as a stand alone unit for deposition of various films113,183. Confined

plasma is extended from the anode region by a positive bias on the chamber with

respect to the anode. The degree of ionisation is controlled by increasing the axial

extraction potential, the electron current and the gas pressure. a-C:H films are

produced from a mix of argon and methane gas, and both Ar and CH4 are introduced

into the system via gas supply vacuum lines.

While it is possible to operate the system on a pure hydrocarbon gas of choice,

in fact this option does not appear attractive as plasma discharge is not stable due to

the poorly controlled ionisation process. Details for the deposition parameters such

as axial current, axial potential, partial pressure, gas mix and their effect onto the

system performance have been studied in detail and reported by Varga5.

The fabricated a-C:H film properties are influenced by many factors including

plasma quality102,184, ion energy and ion density185,186 , plasma – substrate distance

which is related to the mean free path at a given vacuum level, bias potential187,188,

film electrical resistivity189,190, substrate temperature191 and others, and a strong

interdependence between these various factors exist. In the ICP system soft and hard

a-C:H films could be formed under a typical set of experimental parameters:

magnetic mirror potential Vm = 50 V and current Im = 200 mA, axial potential Va =

155 V rms at 50 Hz and current Ia = 220 mA, gas flow rate 9 std. cm3/min; pressure

CH4 gas, PCH4 = 3.5 x 10-4 Torr and Ar gas, PAr = 3.5 x 10-4 Torr. The magnitude of

the magnetic field at the centre of the magnetic mirror is approximately 22 G, which

is almost 8% of the peak field value 5. The substrate bias Vs = 500 V DC and + 300 V

37

rms at 50 Hz, whereas the positive ion current reaching the substrate is Is = 2.2 mA.

The AC component of the substrate potential is set out of phase to the axial potential

Vs and it can be adjusted to give either a net positive or net negative substrate current.

Fig. 22. Open plasma source reactor (courtesy of Laserdyne Pty Ltd).

a-C:H films could be formed at relatively high deposition rates, however low

deposition rates of about 1 – 2 micron/hour are preferred for quality films. Since CH4

is used as the main ion source gas, it is impossible to avoid hydrogen inclusion in the

growing film. However, the absorption of hydrogen ions has been eliminated by

reducing the axial potential to below 160 V 5.

38

3.1.1 a-C:H experimental arrangements

The following experimental arrangements were used to fabricate a-C:H

samples for the project. The films were deposited onto optical grade Si <100>

substrate. Prior to deposition, all substrates were sputter cleaned in pure Ar plasma at

negative bias voltage of Vs = 300 V for 5 min. A mixture of Ar and CH4 was used in

the ratio of 45% Ar and 55% CH4 for the plasma source. The variation of the

negative DC bias voltage was selected as a main deposition variable for fabrication

of a-C:H films since it was known from the literature to be one of key factors

affecting the formation of DLC films in the system similar to the ICP144,187,188,192,193.

The reported negative bias range is within 50 – 900 V DC; soft polymeric a-C:H

films are usually produced using the low values of 50 – 200 V and ta-C:H films are

fabricated using the high bias range of above 700 V. In our experiments hard a-C:H

films will be fabricated using the medium bias which will produce samples with a

narrow range of the sp3 bonding distribution. Other ICP system parameters were: the

magnetron current during deposition Im = 200 mA and the discharge potential Vm =

90 V, the axial current Va = 150 A and the axial potential Ia = 150 V rms, at 50 Hz

frequency. The vacuum level was maintained at 5.5×10−4 Torr during the entire

deposition process. The temperature on the surface of the film during deposition was

400 ± 10 K monitored ex situ using thermocouple measurements. A deposition time

of 4 h (±2 min) was used for all samples fabricated in this study. The flow of gas

mixture was maintained constant at the rate of 8 std. cm3/min in order to keep the

deposition procedure stable. The substrate bias voltage, Vs, was applied onto the

rotating aluminium substrate and was varied between −250 V and −400 V in 50 V

increments for deposition of samples at four substrate bias settings.

3.2 Growth of hydrogen free DLC using the RIBSD system

Conventional IBSD84-86 technique usually employs either a single ion source to

bombard a target, as in Fig. 23, or two ion sources are used, where one ion beam is

39

used to sputter a target and the other to bombard a growing film, Fig. 17. When an

additional beam is used in IBSD deposition, extra energy is supplied to the growing

film that promotes favourable morphological changes in the forming films87,88.

Fig. 23. Ion beam sputter deposition IBSD (single beam).

Fig. 24. Schematics of RIBSD, where a single ion beam is used for target sputtering

and concurrent substrate bombardment.

The outline for the IBSD and the RIBSD techniques was presented in Sections

1.2.1.2 and 2.1.3, and here we summarise several key questions regarding DLC

growth investigation using the RIBSD.

These are:

- How does the incidence angle of the ion beam to the target, denoted αt affects

the growth of DLC films. The incidence angles investigated are 15, 30 and

45°, Fig. 25.

- How the incidence angle of the ion beam on the substrate, denoted αs affect

the growth of DLC films (see Fig. 25). The values of αs investigated is 0

(parallel to the ion beam axis) and 10°.

40

- How does the energy of the incoming ion beam affect the formation of the

DLC film and promote the sp3 bonding. It is expected that the higher energy

of incoming ions will result in higher degree of the sp3 fraction in fabricated

films. Ion energies in the range of 0.1 eV to 1.2 keV (efficiency range for the

40 mm Kaufmann ion gun) are investigated.

- How ion species of different size and atomic mass affect the formation of

DLC, keeping all other deposition parameters fixed (i.e. same sputtering

geometry, ion beam energy). Ar and Xe ions are used in this investigation

(99.999 % purity gases).

Fig. 25. Schematics of relative target and substrate positions to the incident ion

beam flux.

3.2.1 Monte Carlo simulations of Ar and Xe ions interacting

with a target

To optimise the sputtering geometry for the RIBSD experiments Monte Carlo

(MC) calculations were performed where sputtering ions (Ar and Xe) were used to

model the possible outcomes of the experiments. It is known that quality DLC are

formed when the C ions have the energy of approximately 100 eV. At the energy

range chosen for the RIBSD experiments, (0.1 - 1.2 keV) for bombarding with Ar or

Xe ions, the probability of obtaining 100 eV per C ion that has been ejected from the

graphite target due to sputtering is nil. Our hypothesis is that, if the flux of C ions

41

with the energy range of 50 – 70 eV per C ion is obtained from sputtering of the

target, the additional energy will be provided by the impinging ion beam that will

elevate the total energy required to promote the sp3 formation to 100 eV per C ion. It

is believed that the impinging ion beam will also cause the secondary ionisation of

the plasma flux that is formed and confined between the target surface and the

substrate. This flux will contain a substantial amount of backscattered (BS) ions, and

the amount of BS ions will depend greatly on the αt angle. The higher the volume of

BS ions in the confined plasma flux the higher will be the plasma density due to the

constant supply of energetic sputtered ions. The MC calculations were aimed to

obtain an optimum angle when the energy of sputtered C ions to approach the highest

value for the selected Ar or Xe ion bombardment energies.

The sputtering process can be described as follows 6. During sputtering surface

atoms are removed from the target by creating recoil cascades that come back out of

the target, and which give surface atoms enough energy so that they are driven away

from the target. When a cascade gives the target atom the energy that is greater than

the surface binding energy, Eb (see Eq. (1)) of that target, the atom may be sputtered 6. However, to actually be sputtered, the energy of the atom normal to the surface

must still be above the surface binding energy when it crosses the plane of the

surface. The sputtering of a surface is described by a ‘sputtering yield’, Ys, which is

defined as the mean number of sputtered target atoms per incident ion 6. MC

sputtering calculations were performed by using TRIM 6 software. A total of 100.000

ions were used during calculations and the thickness of the HOPG target was set at

50 nm. The 50 nm thickness was chosen to reduce the resources needed for the MC

calculations as a supercomputer power is required for targets with over 1 micron

thickness (apart from saving time the 50 nm choice makes no difference to the final

sputtering results obtained). Simulation were performed with ion energies from 0.1 to

1.5 keV with 0.25 keV increments, and for angle of incidence from 85° (5° to

normal) to 50° (40° to normal).

The percentage of BS ions as a function of the ion energy and the incidence

angles for Ar and Xe bombarded ions is shown on Fig. 26 and Fig. 27 (the shaded

area shows a limit of effective operation for the ion gun). From Figs. 26 and 27 it is

evident that the percentage of total BS ions increase inversely proportion to the ion’s

angle of incidence to the target.

42

0

10

20

30

40

50

60

0 0.25 0.5 0.75 1 1.25 1.5Ar ion energy, keV

% b

acks

catte

red

ions

85 deg 80 deg 70 deg 60 deg 50 deg

Fig. 26. Percentage of BS ions as a function of the ion energy and the incidence

angle. Ar bombardment of HOPG target. Inset shows the angle definition used for

clarity.

0

10

20

30

40

50

60

0 0.25 0.5 0.75 1 1.25 1.5Xe ion energy, KeV

% b

acks

catte

red

ions

85 deg 80 deg 70 deg 60 deg 50 deg

Fig. 27. Percentage of BS ions as a function of the ion energy and the incidence

angle. Xe bombardment of HOPG target.

For Ar sputtering a HOPG target this relationship is clear. However, for Xe sputtering

number of BS ions appear to follow a log scale at low ion energies from 0.25 keV to

θ target

ions

43

approximately 1.0 keV at low ion incidence angles to the target. It appears that for

the range of Ar and Xe ion energies sought for the RIBSD experiments the optimal

angle αt will be approximately 70° or less (optimal utilisation of the total ion beam

flux and thus achieving a better sputtering efficiency). Higher amount of BS ions that

are ejected from the target surface at higher incidence angles (more than 70° when

sputtering with Ar ions and more than 80° for Xe ions) indicate that Ar is better

suited for sputtering, as the high density BS cascade can be obtained at a relatively

low incident angle of 70°. 10 % of BS ions will constitute to the total plasma flux

that is formed between the target and the substrate.

The energy of sputtered carbon atoms as a function of the incident ion beam

energy for Ar and Xe ions at varying incident angles is shown in Figs. 28 and 29.

These figures illustrate that at ion bombardment energies of 1.25 eV or less there is

no flux of C ions produced with the energy range of ~ 100 eV, but much less than a

half that value (i.e. 40 – 50 eV). In order to produce a flux of C ions with energies of

~100 eV for example, the Ar and Xe ion bombardment energies should be well above

the calculated 1.5 keV range, but 9 – 15 keV 6. The stable operating threshold for a

40 mm Kaufman ion source is within 0.3 – 1.0 keV range. This operating range

corresponds to the C flux yield of ~30 eV/atom at 70° incidence. Therefore, as

mentioned before, the impinging ion beam is thought to be a solution delivering the

additional energy to the forming DLC. Figs. 28 and 29 show that the sputtering yield

is significantly higher at high incident ion angles of 80° and 85°, however the

incident flux is reduced by the high angle of incidence and is only a half of what it is

at 70° and 65°. This observation correlated with the BS data from calculations shown

in Figs. 26, 27 and allows us to assume that the favourable angle of incidence would

be αt ≈ 70° (30° to normal).

Figs. 30 and 31 illustrate calculations of the ‘bulk yield’, or the number of C

atoms sputtered per a single incoming ion at a given energy. Careful analysis of Figs.

30 and 31 indicate that at the ion energies up to 1.25 keV the yield of atom to ion is

very similar to that at the ion incident angles of 85°, 80° and 70°.

44

0

10

20

30

40

50

60

70

80

0 0.25 0.5 0.75 1 1.25 1.5

Ar ion energy, KeV

Yiel

d eV

/ato

m

85 deg 80 deg 70 deg 60 deg 50 deg

Fig. 28. The relationship between the energy of the sputtered C atoms, the incoming

ion bombardment energy and the angle of ion incidence. Ar bombardment of HOPG

target.

0

10

20

30

40

50

60

70

80

0 0.25 0.5 0.75 1 1.25 1.5Xe ion energy, KeV

Yie

ld e

V/at

om

85 deg 80 deg 70 deg 60 deg 50 deg

Fig. 29. The relationship between the energy of the sputtered C atoms, the incoming

ion bombardment energy and the angle of ion incidence. Xe bombardment of HOPG

target.

45

0

1

2

3

4

5

6

7

8

0 0.25 0.5 0.75 1 1.25 1.5

Ar ion energy, KeV

Yiel

d at

om/io

n

85 deg 80 deg 70 deg 60 deg 50 deg

Fig. 30. Number of C atoms ejected per single Ar ion at a given incident angle.

0

1

2

3

4

5

6

7

8

0 0.25 0.5 0.75 1 1.25 1.5

Xe ion energy, KeV

Yie

ld a

tom

/ion

85 deg 80 deg 70 deg 60 deg 50 deg

Fig. 31. Number of C atoms ejected per single Xe ion at a given incident angle.

There is considerable reduction of the yield at the incident ion angles of less than 70°

for Ar and Xe. Again, this observation allows us to presume that the favourable αt ion

46

beam incident angle for the target sputtering will be approximately 70° (30° to

normal).

3.2.2 The RIBSD experimental arrangements

Hydrogen free amorphous carbon (a-C) with predominantly sp2 bonding

arrangement and DLC films were fabricated in the RIBSD experiments using

selected target, αt and substrate αs angles, the ions types and the ion energies (Section

3.2). The details for the sputtering geometry expressed as αt : αs, the ion types used

and their respective ion energies for the RIBSD experiments are summarised in

Table 2.

Table 2

The ion types used during the RIBSD experiments, their respective energies and the

angles αt and αs (expressed as αt : αs (in °)).

Ion

energy

keV

Target and substrate to the ion beam axis sputtering angles, αt : αs

15° : 0°

15° : 10°

30° : 0°

30° : 10°

45° : 0°

45° : 10

0.2 --- --- Ar, Xe --- --- ---

0.4 --- Ar Ar, Xe --- --- ---

0.6 Ar Ar Ar, Xe Ar, Xe Ar Ar

0.8 Ar --- Ar, Xe --- Ar ---

1.0 --- Ar Ar, Xe Ar, Xe --- Ar, Xe

1.2 --- --- Ar, Xe Ar, Xe --- ---

47

The actual RIBSD experimental set up is shown in Fig. 32 which illustrates a

Kaufmann gun and a target/substrate holder; the RIBSD in operation is shown in Fig.

33.

Fig. 32. The RIBSD experimental set up (courtesy of Laserdyne Pty Ltd). Kaufmann

ion source is shown on the left and the target/substrate holder is on the right.

Fig. 33. The RIBSD in operation. Argon plasma discharge is visible between the

body of the ion gun and the target/substrate holder.

48

The RIBSD deposition was carried out on a CTP-700 high vacuum deposition

system (Laserdyne Pty Ltd) fitted with a Kaufman-type ion source with a convex ion

grids of 40 mm in diameter. HOPG target of 65 mm x 85 mm and 7 mm thick was

used as the target material. Immediately prior to the sputtering experiments the

HOPG piece was annealed in vacuum at 10-5 Torr pressure at the temperature of 300

K for 30 min. The substrate plate 65 mm x 85 mm was made out of mild steel and

was electrically connected with the target thus minimising the surface charge

accumulation during the deposition. The distance from the grid of the ion beam to the

joint line where the target and the substrate plate meet was approximately 165 mm.

The alignment of the Kaufmann gun to the target-substrate holder relative to the base

of the chamber was performed using an alignment laser. During the RIBSD

experiments the ion beam voltage was varied from 0.2 keV to 1.2 keV with 0.2 keV

increments and a-C and DLC films were deposited onto optical grade Si <100>

substrates. The substrates were positioned on top of the substrate plate. The operation

variables during the RIBSD experiments are summarised in Table 3.

Table 3

Operation variables during RIBSD experiments.

Base vacuum 2 x 10-5 Torr

Working pressure 4.1 - 4.2 x 10-4 Torr

Ion beam 0.2 - 1.2 keV at 10 mA

Accelerator 190 V at 1.0 mA

Discharge 60 V at 0.6 A

Filament 12.5 - 14.0 A

Neutraliser 2 mA, DC

Deposition time 30 ± 1 min

49

Chapter 4

4 Analytical methods used to study the fabricated

DLC films

To begin with, we present the band structure of amorphous carbons, list and

discuss selected analytical methods that are currently used by the carbon community

to obtain reliable information about DLC microstructure. The methods used to

analyse the fabricated DLC films were: nanoindentation measurements, Raman (vis -

UV), FT-IR, NI-R and X-ray photoelectron spectroscopy and scanning probe

spectroscopy (SPS) are also presented and discussed.

4.1 Approach to DLC structure - properties

characterisation

The discussion about the hybridised states of carbon, the σ and π bonds was already

presented in Chapter 1 and 2, and Fig. 1. The vibrational density of states (VDOS)

band diagram1, Fig. 34 for carbon shows that the σ bonds (sp3 atoms) of all carbon

sites form occupied σ states in the valence band and empty σ* states in the

conduction band separated by a wide σ – σ* gap. The π bonds of sp2 and sp1 sites

form filled π and empty π* states, with a much narrower π – π* gap. This simple

model for carbon atomic structure was developed some years ago, and it was based

on properties of σ and π bonds194,195. It has been argued that maximising the π

bonding energy results in sp2 sites forming π bonded clusters within the sp3 bonded

matrix. The argument was supported by the cluster size observation which

determines the actual band gap. Later it was found out that the cluster model grossly

50

overestimates the actual size of the cluster80,196,197, however, the initial hypothesise

was correct1.

Fig. 34. Schematic VDOS of a carbon showing σ and π states. From Ref. 1,198.

The ternary phase diagram for the C-H system shown in Fig. 2 emphasises that

there are two key parameters which determine the structure and properties for

hydrogenated DLC: the fraction of sp3 bonded sites and hydrogen content. The

ordering of sp2 sites is a third significant factor and it is important for the electronic

properties.

Table 4

Comparison of characterisation methods available for DLC analysis, their advantages

and availability.

Methods Comments Availability, usage

Nuclear magnetic

resonance

Indirect, large sample required, C13,

dephasing Available*

X-ray diffraction Time consuming Available*

XPS spectroscopy Indirect, small peak shifts Available, used

IR, N-IR spectroscopy Only sites bonded to hydrogen Available, used

Raman spectroscopy

(visible) Indirect, sp3 sites invisible Available, used

Raman spectroscopy

(UV)

Indirect, sp3 and sp2 visible, 244 nm

preferred Available, used

*not used due to the difficulty of separating the DLC film from the substrate

51

Various characterisation methods have been developed to determine these three

important structural parameters: the amount of sp3 and sp2 phases and the hydrogen

content. Near-edge X-ray-absorption spectroscopy (NEXAS)199, high-energy

electron energy loss spectroscopy (HEELS or/and EELS)200-203 and X-ray

reflectivity204-206 are currently considered the most reliable techniques to determine

the sp3 fraction in DLC. One should distinguish between the methods suitable for

detailed analysis such as diffraction, and more routine methods for repeated

structural monitoring focused primarily on determination of the sp3 and hydrogen

content. The effectiveness and disadvantages of various available methods are

summarised in Table 4. From Table 4 we can see that only indirect characterisation

methods are available, and used during throughout the course of the project to

determine the sp3 and hydrogen content, these were the XPS, FT-IR, N-IR and

Raman spectroscopy.

4.2 Nanoindentation measurements

The mechanical properties of DLC films are of great importance since the films

are known to be used as protective coatings in aggressive environments60,140,207-212

and for tribological applications44,213-216. DLC films have an advantage over

polycrystalline diamond dues to an absence of grain boundaries and by providing a

good coverage on large areas on most of industrial substrates.

Mostly the mechanical properties of thin films are measured by using a nano-

indenter44,217-220 where a load is applied to a small indenter (usually diamond) and the

depth of penetration beneath the specimen surface is measured. Data for applied

force and depth is then collected during loading to the prescribed maximum load and

also unloading back to the zero load221, Fig. 35. The data could be obtained either by

using the single-point load-unload method or the multiple-point load-unload method

(Oliver and Pharr222). The only difference between the multiple-point load-unload

method and the single-point unload method is the number of unload data points used

to determine the elastic modulus and hardness values221.

52

Fig. 35. Multiple-point unload method uses the slope of the tangent to the initial

unloading to determine hp. Single-point unloading is faster, and hence it is less

affected by the thermal drift, but relies on a single data point in the unloading portion

of the test cycle (From Oliver and Pharr221).

The multiple-point load-unload method uses several data points on the unload part of

the test and fits a curve to these data. The slope of this curve is then extrapolated to

the depth axis and various corrections are applied to this intercept that determine the

plastic depth hp. The slope of the line is used to determine the combined modulus of

elasticity, E*, using the derivative of the Hertz elastic equations221

i

2i

s

2s

Ev1

E1

*E1 −

+−

=ν

(14)

where, Es and Ei are the Young’s modulus and νs and νi are Poisson ratios of the

sample and the indenter respectively.

Then, the combined modulus E* can be expressed as

21

AdhdP*E π

= (15)

53

where, dP/dh is the slope of the loading curve, dh is the penetration depth, dP is the

loading force, Fig. 35, and A is the area of an indenter 221.

For the Berkovich indenter, the relationship between the loading force and the

penetration depth is221,223,

( ) ( )⎥⎦

⎤⎢⎣

⎡⎟⎠⎞

⎜⎝⎛ −

+=−

*21

2

E2H22tanH33Ph π

ππα

(16)

where, h is the penetration depth, P is the loading force, H is hardness, tan α is slope

of the unloading curve and E* is the combined Young’s modulus (Eq. 15). The

penetration depth h, (Eq. 16) is calculated by taking into an account both the elastic

and the plastic deformation modes and it is calculated based on an assumption that

there is an elastic-plastic contact throughout the measurement cycle221.

Finally, hardness H, is defined as

APH =

(17)

where, P is the load and A is the area of the indentation measured from the

plastic deformation depth.

It is well known that in order to obtain accurate measurement results, the

indentation depth must be limited to a fraction of order 15 – 20 % of the total film

thickness222,224,225. The compliance with this rule is important for both the hard and

the soft films.

4.2.1 Instrumental settings UMIS

The UMIS 2000 Ultra Micro-Indentation System221,223 was used to measure

hardness and elastic modulus of hydrogenated DLC films. The maximum indentation

load used was 6 mN. The indentation load measurements started from a low load of 2

54

µN, that was used to identify the surface contact, and 0.01 mN increments were used

for load and unload cycles. The analysis of the load-displacement curves were based

on Oliver-Pharr method222 using the Berkovich indenter and hardness H, and

Young’s modulus E, were calculated according to this method. Since H and E

parameters are directly influenced by the surface area of the indenter calibrating

calculations were performed to account for a non-ideal shape of a real indenter prior

to the measurements. Corrective parameters accounting for a non-ideal area function

were obtained from a test on a standard fused silica sample, with H = 9.00 GPa, E =

72.5 GPa and Poisson ratio ν, of 0.17. All mechanical properties calculations were

performed for 100% load - unload vs. displacement curves. Experimentally, each of

the load/unload cycle was performed with 20 unload increments. The dwell time

during a single incremental stop was 0.5 seconds. The distance between each of the

penetration point was 20 µm and the total examined area for each of a-C:H sample

was 225 mm2. Corrections for the thermal drift, initial penetration depth, instrument

compliance, and a non-ideal indenter geometry were used, as specified in the UMIS

User Manual223. The shape factor Cs, accounting for non axis-symmetric nature of

the indenter was 1.034. The intercept factor Ce, used for corrections in multiple point

unload method was 0.72. The thermal drift correction was calculated as221,223

tThh̀ d−= (18)

where, h` is the corrected depth, h is the measured depth, Td is thermal drift nm/sec,

and t is time.

Instrument compliance was determined as221,223

PChh̀ f−= (19)

Where, h`, again is the corrected depth, h is the measured depth, Cf is deflection of

the load frame (Cf = 1 nm/mN), and P is the applied load.

The hardness of the Si substrate on which the a-C:H samples were deposited

onto was measured prior to conducting all measurements and was found be 13.2 ±

0.1 GPa.

55

4.3 X-ray photoelectron spectroscopy and obtaining sp2/sp3

ratio from the core level C1s peak

X-ray photoelectron spectroscopy (XPS) is a useful tool for investigating the

local binding configuration of various materials and has been used widely in

materials science226-229. Due to the localized nature of the core-level electrons, the

XPS measurements are predominantly sensitive to their localised potential and thus

information about a local chemical bonding environment for specific atoms can be

obtained. XPS spectra of amorphous carbons usually show a broad C1s peak

composed of the binding energies of natural diamond and graphite230-232. The shape

of the C1s peak is found to depend mainly on the atomic density of the amorphous

samples. This observation is generally interpreted through the analysis of C1s line

consisting of two components associated with the sp2 and sp3 hybridized carbon

atoms. Relative intensities of the two contributions are considered to provide a

reliable measure of the sp2/sp3 hybridization ratio30,232-237. The individual binding

energy (BE) for the sp3 and sp2 phase (eV) arises from a chemical shift value that

results of an effective electron charge transfer 238. The crystalline diamond C1s peak

is positioned at ~ 285.3 eV, Fig. 36 and is positioned approximately 0.8 - 0.9 eV

higher than of the graphite at ~ 284.5 eV239,240.

Fig. 36. The C1s diamond spectrum. From Ref. 235.

56

This separation is attributed to the sp2 atoms of having a shorter bond length of

~1.41 Å and are a subject to a deeper x-ray ionisation potential as compared to the

sp3 atoms with the bond length of ~1.54 Å241. The shift between sp2 and sp3 phases is

also attributed to the difference between the lattice constant of the sp3

diamond116,242,243 that is ~3.56 Å and sp2 graphite244-246 that is ~2.46/6.71 Å, therefore

more energy is required to remove an electron from the sp3 site. The lattice constant

is assumed be different for hydrogenated and hydrogen free amorphous carbons. The

full width at half maximum (FWHM) of diamond C1s peak is usually 1.0 – 2.2 eV

and it is almost twice as wide as the FWHM of natural graphite that is only 0.6 – 1.0

eV247-249. This discrepancy between FWHMs for diamond and graphite results in the

FWHM for amorphous carbons of being wider than a sole FWHM of the sp3

diamond, since it is a superposition of two lines corresponding to the sp2 and the sp3

hybridised states241,248, as illustrated in Fig. 37.

Fig. 37. Measured C 1s photoelectron spectra of a a-C film. The Shirley background

and the sp2 and sp3 components resulting from the fit are also shown 241.

The information about the nature of atomic bonding in DLC is obtained from

C1s spectra the following way. The C1s spectra is corrected to account for the Shirley

background250 that is believed to be due to extrinsic electron energy variations.

However, very recently Vegh251 reported that any background correction in XPS is a

questionable undertaking due to unintentional subtraction of important intrinsic

contributions built in any subtraction model used. Still, in the absence of defined

alternatives most researchers (including in this work) subtract a background using the

57

Shirley line-function. Then, the comparative BE spectral functions of graphite and

natural diamond234,247 are used to calculate the sp3/sp2 ratio based on the area ratios

of respective hybridised functions233,248 in compliance with the binding energy

difference ∆BE between the sp2 and sp3 constituents when fitting the spectra.

∆BE is calculated as

∆BE = BEsp3 - BEsp2 (20)

where BEsp3 and BEsp2 are the binding energies of the respective hybridised states,

The value of ∆BE is found in the 0.80 - 0.95 eV range for both hydrogenated

and hydrogen free DLC containing any amount of sp2 and sp3 bonding. This range

was obtained experimentally by referencing the XPS C1s peak analysis results against

the results obtained by EELS166,233,234,239,252 that is known to provide reliable sp3

probing. The XPS C1s spectra analysis provides incoherent results4,30,253 when ∆BE is

used outside the established range.

Usually the C1s core level peak is fitted with either pure Gaussian, Lorentzian,

the combination of Gaussian and Lorentzian232 or Doniach – Sunjic241,254 line-shape

functions. The Gaussian line-shape accounts for the instrumental broadening, the

chemical disorder and the phonon interaction, while the Lorentzian function accounts

for the finite lifetime of the core hole in the photo-ionization process. The use of

unmodified Gaussian or Lorentzian line-shapes is perhaps, not appropriate as there

are parameters that affect the shape of the intrinsic function (see Section 4.4.1 for

Raman). The most appropriate will be the use of a combination function, for example

of Gaussian and Lorentzian functions, which interact at various amounts thus

ensuring the instrumental broadening and photo-ionisation processes are also

accounted for. In the analysis of C1s spectra the authors used Pearson VII line-shape

function (see Section 4.4.1 for full details) that is composed of both the Gaussian and

Lorentzian line-shape functions.

58

4.3.1 Instrumental settings XPS

The XPS C1s measurements were performed using Kratos Ultra photoelectron

spectroscope with monochromated and non-monochromated Al Kα 1486.6 eV X-ray

sources. The chamber vacuum level was maintained below 2 x 10-11 Torr all during

measurements. The C1s core-level spectra were obtained with x-ray constant pass

energy at 50 eV and the total experimental resolution of 0.25 ± 0.01 eV; the number

of sweeps were 4. The spectrometer was calibrated by peak referencing of Au 4f7/2

(BE of 84.0 eV) with respect to the Fermi level. The measured photoelectron

intensity was corrected for the analyser transmission function proportional to E-0.75,

where E here is the photoelectron kinetic energy. Lateral resolution for XPS

measurements was approximately 800 μm.

4.4 Multi-wavelength Raman spectroscopy

Raman spectroscopy is a popular, non-destructive tool for structural

characterisation of amorphous carbons255-262. The measurements are traditionally

carried out in the blue–green spectral region (488 – 514.5 nm), however, multi-

wavelength (MW) Raman studies are now gaining popularity. Recently there has

been a considerable improvement in the field of MW Raman spectroscopy of carbon

systems. In particular, the appreciation of the strict correlation of the Raman process

with the electronic properties of carbon systems is a major driving force to further

develop all the possibilities of this versatile technique 1. MW Raman has recently

been used to investigate the origins of the peaks at ~ 1150 cm−1 and ~ 1450 cm−1 in

nanocrystalline diamond263, and it has also been used to study the carbon nitrides264

and an isolated carbon nanotube265-268. The Raman spectra of diamond, graphite and

some disordered amorphous carbons are compared in Fig. 38. All carbonaceous

materials in the visible Raman excitation spectrum display common characteristic

peaks in the first order region, up to 2000 cm-1 region 1.

59

Fig. 38. Raman spectra of carbonaceous materials. From Ref.1.

In this region the E2g vibrational mode of graphite of D46h symmetry is assigned

to vibrations of atoms in a poly-aromatic arrangement. The first mode, the E2g1, is a

low frequency shear mode centred at around 40 cm-1. The E2g1 mode corresponds to

relative vibrations of the atoms in the plane perpendicular to the aromatic layers. It is

seldom studied due to the difficulty of separating out the Rayleigh background.

Diamond has a single Raman active mode at 1332 cm-1, which is a zone centre mode

of T2g symmetry. Single crystal graphite has a single Raman active mode, the G

mode, for ‘graphitic’, centred at about 1560 cm-1. Disordered graphite has a second

mode, denoted the D mode, for ‘disorder’, of A1g symmetry centred at approximately

1360 cm-1. The G and D peaks are features of the sp2 hybridised fraction only. The G

corresponds to sp2 bond stretching vibrations in both aromatic chains and olefinic

rings. The D peak is due to aromatic ‘breathing’ vibrations of sp2 rings 1. An unusual

and significant fact is that the Raman spectra of most disordered amorphous carbons

remain dominated by these two G and D modes of graphite, even when the

amorphous carbons do not have any specific graphitic ordering 260.

Raman is light scattering by the change in polarisability χ due to a lattice

vibration1,157,

60

)q,k(Qdqd)k( 0χχχ += (21)

where χ is the polarisability at a wavevector k, and Q is the amplitude of a photon of

wavevector q.

This change in polarisability causes a non-elastic scattering of an incident

photon with inherent (ω, k) into the scattered photon (ω`, k`), where ω is the photon

frequency. The Raman cross section can be expressed as

2

ddkC ⎟

⎠⎞

⎜⎝⎛=

ωχ (22)

The polarisation can occur by excitation of the electronic ground state into virtual

states at energy E, or into real states at E1.

In an amorphous material like DLC, there is a complete loss of periodicity and

a breakdown of the k selection rule of optical and phonon transitions1. In this case,

the IR and Raman spectra of an amorphous network correspond to the VDOS, G(ω),

that is weighted by an appropriate matrix element C(ω). This is the Shuker –

Gammon formula for the Raman spectrum1,269.

)(G)(C1)(n

)(I ωωω

ωω

+= (23)

The Raman and IR spectra are usually relatively smooth and resemble each

other, and this resemblance is noted for a-C:H films in Sections 4.4 and 4.5, however

close spectral matching is only true for certain films types while most striking

similarities were only observed for Raman excitations for the visible range. The

reason is that the Raman spectrum in visible excitation is dominated by G and D

modes that are originated by scattering from the sp2 sites only. The π states are of

lower energy than the σ states, therefore they are much more polarisable1,194. This

gives the sp2 sites a cross-section of 50 - 230 times larger than the sp3 sites270,271, that

is why the sp2 sites dominate the Raman spectra of even for high quality ta-C films

61

which only have a residual 10 – 15 % sp2 content. This is clearly seen in Fig. 39 (a).

It is now known that Raman peaks disperse with varying excitation energy in

amorphous carbons, Fig. 39 (a) and (b).

Fig. 39. MW Raman spectra of (a) ta-C, (b) ta-C:H, (c) sputtered a-C and (d)

polymeric a-C:H. The peaks’ trends are indicated. From Ref. 260,262.

Raman scattering in UV excitation reveals sp3 sites as indicated by the

appearance of T peak at approximately 1050 cm-1 1,262,272,273. The sp3 fraction is

undetectable in visible excitation wavelengths since low energy visible light does not

excite the higher lying σ states (Fig. 39) and only allows probing across the π – π*

gap. Probing of the σ – σ* gap (Fig. 39) requires much higher energy than the

visible spectrum is able to provide and UV Raman at wavelength of 244 nm (~5.1

eV) provides equal excitation for both the sp2 and sp3 sites271,274.

62

The analysis of DLC microstructure when using Raman in the visible range

(only sp2 bonding) is customary performed by monitoring the relative peak intensity

I(D)/I(G) ratio, the G peak position and the FWHM of the G peak, FWHMG. For the

UV excitation: the relative intensity of the T peak over the G peak, I(T)/I(G) and the

T peak position are additionally used.

There is a spectral dispersion effect, Fig. 39, due to photon confinement in π

and σ VDOS that is described using Eq. 21. Example of the G and T peaks

dispersion trends are indicated in Fig. 40. Fig. 40 (A), shows the variation of the G

peak position with laser excitation wavelength and energy. The G peak does not

disperse in graphite itself nor it does in nanocrystalline (nc)-graphite or glassy

carbon275,276.

Fig. 40 (A, B1, B2). A: The dispersion of the G peak vs. excitation

wavelength/energy for a series of template samples. B1 and B2: I(T)/I(G) and T peak

positions vs. sp3 fraction for non-hydrogenated carbon films. From Ref. 260,262.

The G peak only disperses in more disordered carbon systems, and the dispersion is

proportional to the degree of bonding disorder1. The behaviour of the G peak in

disordered graphite is very different from amorphous carbons, even though the G

peak positions might coincide. The G peak in graphite cannot disperse because it is

the Raman-active phonon mode, the G mode, of the crystal. In nc-graphite, the G

peak shifts slightly upwards at fixed excitation energy due to phonon confinement,

63

but it cannot disperse with varying excitation energy owning to the VDOS rule1. The

dispersion occurs only in a more disordered carbon system owning to different local

band gaps and different phonon modes. The G peak dispersion separates materials

into two types, in those with only sp2 rings, the G peak dispersion saturates at a

maximum of ~ 1600 cm−1 that is the G position in nc-graphite1. In those materials

also containing the sp2 chains (ta-C, ta-C:H), the G peak continues to rise past the

1600 cm−1 barrier and can reach a maximum of 1690 cm−1 at 229 nm excitation in ta-

C. This high G peak position can only be due to short, strained C=C bonded chains1.

Fig. 40 (B1 and B2) gives some empirical relations between the I(T)/I(G) ratio,

the T peak position and the sp3 content260,277,278. Notably, the variation of I(T)/I(G)

with the sp3 content is quite non-linear for 60 – 90% sp3 contents, as shown in Fig.

40 (B2). On the other hand, as the sp3 content falls, the VDOS peak at 1060 cm−1

shifts upwards to that of a sp2 network at 1400 cm−1. Alternatively, the changes could

be represented as a reduction of the T peak at 1060 cm−1 and the rise of a peak at

approximately 1400 cm−1, a D-like peak. Thus, as the sp2 content of ta-C rises, the T

peak intensity corresponding to the C-C sp3 VDOS is reduced, with a corresponding

increase of a D peak. A complication is that the D peak intensity depends not only on

the sp2 fraction, but fundamentally on its order. If the sp2 sites have graphitic order,

the D peak is absent in UV, if the sp2 sites are in chains the D peak is also absent,

and only if the sp2 sites are in disordered rings does a residual D peak survive in

UV260. This can explain the range of I(T)/I(G) values seen for high sp3 content ta-C1.

The increase of sp2 content and clustering both tend to reduce the T peak

intensity relative to the G peak. However, the T peak disappears only for large sp2

contents. Thus, the clustering effect reduces the direct correlation between the T

intensity and the sp3 content. Nevertheless, it is still possible to distinguish high sp3

contents from low sp3, unlike in the visible Raman spectra. Indeed, a T peak at

approximately 1060 cm−1 and an I(T)/I(G) ratio of approximately 0.40 – 0.42 in

hydrogen free DLC samples is a sufficient condition to estimate an sp3 content of

approximately 80 % using the recent works of Ferrari and Robertson259. A I(T)/I(G)

ratio of 0.3 – 0.4 still indicates an sp3 content of 60 – 80%, but sp2 clustering makes

it difficult to give a precise figure. Finally, the I(T)/I(G) of less then 0.2 indicates that

the sp3 content is lower than 20 – 30 %, again using the works of Ferrari and

Robertson259.

64

4.4.1 Lineshape dilemma

Before proceeding with the analysis and discussion of the experimental results,

we need to emphasise on the importance of properly fitting the spectra as it affects

the numerical values. The Gaussian, Lorentzian and skew Lorentzian line shapes are

usually used in deconvolution of Raman spectra. A modified Lorentzian line-shape

function, also otherwise known as the Breit – Wigner – Fano (BWF)1,259 line-shape is

also used. The use of BWF line-shape has been endorsed by the Cambridge

group1,259 and the Prawer group279 and BWF line-shape is becoming very popular

among other researchers273,280. We disagree with the use of BWF for spectral fitting

and argue that it is lacking a scientific basic.

Our argument is presented as follows:

1. Fitting the Raman spectra with the Gaussian line shape could be safely

discarded since the Gaussian fundamentally is a probability function of the normal

distribution that is ideally suitable to describe a statistical distribution of a random

process. The distribution of Raman scattered light from a material surface that is

passed via a spectrophotometer is highly unlikely to be described effectively when

using only this function.

2. The Lorentzian is the natural line shape in spectroscopy. The underlying

physical process described by Lorentzian function is that, after the initial excitation,

the location of the energy level of the exited state is not exactly known and still, there

is some ambiguity about the excitation response which remains. The Raman excited

state is time dependent and it directly influences the FWHM of the Lorentzian line.

When the lifetime of the excited state approaches the transition frequency the

FWHM broadens significantly, and still, the process is described by the Lorentzian

distribution. The Lorentzian is therefore the most appropriate line shape to be used in

spectral deconvolutions, including the fitting of Raman spectra.

3. The BWF1,281 is an asymmetric Lorentzian line shape that is adjusted with the

skewness coefficient, Q, and BWF is expressed as

( )[ ]20

200

/Γωω21

/QΓωω21II

−+

−+=)

⎥⎦⎤

⎢⎣⎡ ⎟

⎠⎞⎜

⎝⎛

ω( (24)

65

Where I0 is the peak intensity, ω – the peak position, Г is the FWHM.

A symmetric Lorentzian corresponds to Q = ∞. Where the skewness of the maximum

of the BWF occurs at

Q0max

2Γ

+= ωω (25)

Therefore, the BWF function could be understood as describing a process the

lifetime of which in the excited state is known, but that is not true.

4. While the inherent line shape in spectroscopy is Lorentzian, it is always

influenced by other line broadening factors such as molecular collisions and the

Doppler effect, thus contributing to creation of an intrinsic line shape. The width of

an intrinsic line shape is also affected by the spectrometer hardware, such as

interactions of a laser source and the instrument response factor. The line shape that

accounts for these and other interactions, broadening factors, etc. is a multifaceted

line shape function known as the Voigt function. The Voigt line shape is a peak that

is based on the combination of Lorentzian and Gaussian functions which interact in

varying amounts, as expressed in Eq. 26.

∫

∫

∞

∞−

∞

∞−

+−

⎥⎦

⎤⎢⎣

⎡−⎟

⎠⎞

⎜⎝⎛ −

+

−

=

22l

2

2

g

02l

20

ydy)yexp(

y

dy)yexp(I

)(

Γ

ΓωωΓ

ωΙ

(26)

where I0 is the peak intensity, ω – the peak position, Гl and Гg are the FWHM for

Gaussian and Lorentzian functions respectively.

The Voigt function is not easily solved analytically, therefore some

approximations are made. One of these approximations is the Pearson VII (P VII),

function that is the broad approximation of the Voigt line shape, Eq. 27.

66

( )M2

1/M00

Γ12ωω21I)I(

−

⎥⎥⎥⎥⎥

⎦

⎤

⎢⎢⎢⎢⎢

⎣

⎡

⎪⎪⎭

⎪⎪⎬

⎫

⎪⎪⎩

⎪⎪⎨

⎧−−+=ω

(27)

where M is the Pearson width.

When the width M approaches 1 or ∞, the Pearson VII (P VII)_ function

resembles Lorentzian or becomes Gaussian correspondingly. The fitting variables

from the P VII function do not correspond to any specific underlying physical

process that originally caused the broadening, as it accounts for the actual presence

of the broadening.

The P VII function was used exclusively throughout the project for the analysis

of Raman spectra of DLC and a-C films.

4.4.2 Rayleigh scattering measurements

It is known that the magnitude of Rayleigh scattered light, IBS, is proportional

to

2

BS *I ⎟⎟⎠

⎞⎜⎜⎝

⎛=

δρδερ (28)

where, ρ is the density and ε is the optical dielectric constant 282.

The density gradient, ρ can also be expressed using thermodynamic constants such

as,

T

T

kdPdnn2

*⎟⎠⎞

⎜⎝⎛

≅δρδερ (29)

67

Where, n is the refractive index and P is the system applied pressure, T is the systems

temperature and kT is the isothermal compressibility. The kT is expressed as

TT d

d1k ⎟⎟⎠

⎞⎜⎜⎝

⎛−=

ρν

ν (30)

These relationships indicate that thermo-mechanical properties of solids do relate to

the Rayleigh light scattering intensity.

In 1999, Voevodin et al283 indicated that there is a strong relationship between

the intensity of the Rayleigh scattered line and density of DLC films examined.

However, the argument was founded on a rather unconventional fitting method of

visible 514 nm Raman spectra with a 1332 cm-1 pure diamond peak together with the

D and G peaks.

In our work we determine the intensity of the Rayleigh line measured at 0 cm-1

using Raman spectrophotometer, and find the relationship between the magnitude of

this line and the relative density of DLC samples. It is believed that high relative

density will correspond to high IBS and the mechanical properties (H and E).

4.4.3 Instrumental settings Raman UV- Vis

Raman spectra in visible excitation wavelengths were collected using

Renishaw Raman spectroscope with 17 mW He/Ne laser excitation source at a

wavelength of 633 nm and Renishaw 87.7 mW SHG laser at a wavelength of 532

nm. The spectra in the visible range were collected at 3 mW laser power in extended

mode with signal collection range of 600 - 3200 cm-1. For Rayleigh scattering

measurements the 532 nm laser was used in the range of -10 – 10 cm-1 in the static

mode with acquisition time of 10 sec.

In the UV range Kimmon 5161R-GS Raman spectroscope with 50 mW He/Ca

laser at a wavelength of 325 nm and 244 nm was used. Since thermal sample damage

is often an issue during the measurements the power of the UV sources was set at or

68

below 1 mW for all measurements, which were performed in dynamic mode by

moving a sample linearly at the speeds of up to 30 μm/sec. Acquisition time for these

measurements was 60 - 480 sec. The grating was centred at 1450 cm-1 for the 244 nm

and an extended range was used for the 325 nm collections. Lateral resolution in all

Raman measurements was approximately 700 μm and the measurements were

collected from an area of ~1 cm2 for each of the sample reported.

4.5 Fourier transform infrared spectroscopy

The probing energy for the Infrared (IR) spectroscopy is defined for

convenience as the near infrared (N-IR) from 0.78 to 2.50 µm; the actual IR (or mid-

infrared M-IR) from 2.50 to 40.0 µm; and the far infrared from 40.0 to 1000 µm284.

The IR absorption for hydrogenated DLC consists of C-H stretching modes at 2800 –

3300 cm-1, whereas C-C modes and C-H bending modes are observed below 2000

cm-1 134,285-292. Fig. 41 shows a summary for a-C:H stretching and bending

modes1,291,292. The middle graph (Fig. 41) illustrates the second derivative that can be

used to separate the peaks since C-H bond bending and rotational modes often

overlay the C-C active modes. In order to confirm the assignment of the IR modes

deuterium substitution is usually used293,294; if the modes are indeed of C-H, their

vibrational frequency positions move by a factor of square root of the mass ratio, that

is 1.4 1. The features of IR active C-H modes follow closely the vibrational

frequency for hydrocarbon molecules. The solid state modes above 1340 cm-1 are

much more localised, and the specific C-H groups can be assigned reliably by

comparison to their values as molecules. Otherwise inelastic neutron scattering can

be used to identify the inactive modes and verify the peak assignments295. IR

ellipsometry is a sensitive way to separate the components of various bands296 and

Ristein et al297 used mass selected thermal fusion to trace the variation of modes

vibrational frequency in IR with hydrogen content and found little variation.

69

Fig. 41. IR spectrum of an a-C:H film. From Ref. 291,292.

FT-IR spectroscopy was used by other workers286,287 together with nuclear

magnetic resonance to study graphite oxide and a-C:H films and good correlation of

the results was obtained. Usually Gaussian line-shape is used to fit sp3 groups and to

derive the sp3 contributions from the IR spectra297.

4.5.1 Instrumental settings IR

Nicolet Nexus Fourier Transform IR spectrometer in 650 – 4000 cm-1 range

was used to collect the spectra. Since studied a-C:H films were deposited onto Si

substrate a transmittance mode was used to collect the signal. The Si absorbance

signal was therefore subtracted to reveal the C:H modes. The IR spectral analysis

was performed by fitting all Gaussian line-shapes. Experimentally, the Eg gap values

(Section 4.6) were derived by extrapolating an absorption curve from the N-IR

spectra.

70

4.6 Band gap (Tauc gap) and surface conduction gap

Electronic and optical properties of amorphous carbons relate to the width of π

– π* gap of the sp2 phase1,10,80,198,298. The band gap is usually derived from the cluster

model that was initially proposed by Roberson et al194. The sp2 clusters in this model

are assumed to be only planar and the band gap can be expressed as

M2Egγ

= (31)

where γ is the V(ppπ) interaction and M is the number of rings in the cluster.

However, due to the limitations of the model (Section 4.2) it is no longer

possible to describe the band gap by a single equation. The theoretical work

performed calculating amorphous carbon networks by Fraunheim et al299,300, Chen

and Roberson et al301,302 and McCulloch at al303 and the experimental work by other

researchers 92,262,304-309 confirmed that the band gap increases with decreasing sp2

content, Fig. 42, and it was also found that the presence of hydrogen has little or no

effect on the band gap since the C-H states lie well away from the band gap region.

The band gap in a crystalline solid is defined as the minimum energy gap

between the occupied and empty states and it can be either direct or indirect. In an

amorphous semiconductor, due to the absence of a strict interconnecting network,

there is no true band gap, and therefore, an arbitrary definition must be used like the

E04 gap defined as the energy at which the optical absorption coefficient, α is

equivalent of 10-4 cm-1, or the Tauc gap, Eg310. The Tauc gap is empirically defined

from the relation

( ) ( )g2/12/1 EhBh −= υυα

(32)

where, hυ is the photon energy, and α is the absorption coefficient and B is a

constant.

71

Fig. 42. Calculated variation of band gap with sp2 fraction. From Ref. 301.

The relationship between Eg and the sp2 content was reported by Babtista et al311 by

analysing empirical results available form the literature17,258,306,312-315, as shown in

Fig. 43.

Fig. 43. Variation of Tauc gap with sp2 fraction 311.

Often, I - V spectroscopy, known as Scanning Tunnelling Spectroscopy (STS)

is used to acquire a normal topographic image of a surface at fixed current and

voltage. Such image is obtained when feedback loop is interrupted at a given position

over a sample surface and the bias voltage (potential between the tip and the surface)

72

is set to a series of voltages, and the tunnelling current is recorded. The voltage is

then returned to a scanning nominal and the feedback loop is turned back on. Each I-

V spectra can be acquired in a few milliseconds for each point on the sample and

such a short acquisition time ensures that there is no appreciable drift in the tip

position. This procedure generates a complete current image at each voltage in

addition to the topographic image316-319. This technique also can be used to obtain an

arbitrary conduction gap for an examined material320,321.

The band gap value that is measured via STS is obtained from the uppermost

layer of the surface. Since this layer is affected by the environment and it usually

only few nanometers thick, such STS measurements are not reliable for obtaining

information about VDOS of the bulk.

4.6.1 Instrumental settings STS

NT-MDT Solver Scanning Probe Microscope was used to obtain the surface

conduction gap, ESC for a-C:H samples. The measurements were conducted at room

ambient atmospheric conditions. Prior to the measurements the sample surface was

cleaned in a mix of ethanol and acetone and dried in nitrogen. The tunnelling current

was approximately 0.3 - 0.8 nA and the sample bias was modulated within 1.5 - 4 V

range. The finite horizontal part on I - V curve at I = 0 nA was used to calculate the

ESC gap of a-C:H samples.

4.7 Scanning electron microscopy

Scanning electron microscopy (SEM) is often used to obtain high resolution

images of a sample surface322-324. Fabricated DLC films were examined using FEI

Quanta 200 Environmental SEM. Prior to examination the samples were cleaned in

73

ethanol and dried in air. The images of the films were obtained at the vacuum level

of approximately 2 x 10-6 Torr in static mode. The filament current was

approximately 2.55 A and the emission current was 105 – 110 µA. The beam current

was 15 – 20 keV for the default spot size of 2 – 3 µm.

The SEM examination revealed that the surface both hydrogenated DLC films

produced using the ICP system and the hydrogen free DLC films fabricated using the

RIBSD technique was featureless and appeared exceptionally smooth at resolution of

300 nm. Since such images do not provide any scientific value we do not include

them in this work. The only exception is the image of an a-C:H sample cross section

where the film is clearly visible relative to the Si substrate. This image is presented in

Section 5.6.

74

75

Chapter 5

5. Characterisation of fabricated a-C:H films

In this Chapter the characterisation of a-C:H films fabricated using the ICP

technique are presented. The full range of characterisation techniques listed in

Chapter 3 is used, providing an opportunity to correlate the bonding characteristics of

the films with the micromechanical and electronic properties of the films.

5.1 Nanoindentation results and discussion for a-C:H samples

The hardness vs. penetration depth diagram for a-C:H samples deposited at

different bias voltage are shown in Fig. 44. These curves have been derived from the

hardness indentation curves for each sample measured using a maximum load of 6

mN and with 20 point load-unload curve as described in Section 4.2 and 4.2.1. The

raw load unload curves for samples at all four bias values used in deposition (-250V,

-300V, -350V and -400V) are shown in Appendix 1. This data was analysed using

Eq. 14 and 15, assuming an indenter hardness of 120 GPa. The thickness of all a-C:H

fabricated films was 9.00 ± 0.02 µm as determined from examining the film cross

section using SEM images see, Section 5.6. The mean penetration hardness Hn, was

approximately 1.60 GPa ± 0.02 GPa. The full penetration depth was found to be

below 6 µm for all 9 µm thick films, as seen in Fig. 44 and Fig. 45. There was a

small variation in hardness, H, between a-C:H samples deposited at -250 V bias,

H=15.2 ± 1.1 GPa and deposited at -400 V bias, H=18.6 ± 1.1 GPa. The Young’s

modulus, E was found to be 174 GPa and 190 GPa for films deposited at -250 V and

-400 V bias respectively.

76

Fig. 44. Hardness as a function of penetration depth for samples deposited at -400 V

and -250 V negative bias. The Si <100> substrate hardness (13.2 GPa) and the mean

surface hardness Hn, are indicated by horizontal lines.

The Young’s modulus vs. penetration depth diagram for the samples as in Fig.43 are

shown in Fig. 45.

Fig. 45. Young’s modulus as a function of penetration depth for samples deposited at

-400 V and -250 V negative bias. The Si <100> substrate E modulus (145 GPa) and

the mean surface Young’s modulus Eio, are shown by horizontal lines.

In Figs. 44 and 45 the initial rising part of the curve at low indentation depths

corresponds to elastic response of the tested sample, while the horizontal section of

the curve around 1 - 2 μm, with a slowly decreasing hardness/Young’s modulus,

77

corresponds to plastic deformation. All H and E values were determined at a

penetration depth that is less than of 1.8 µm in compliance with 15 - 20% penetration

depth rule 224,225. Beyond 2μm, exceeding this allowable penetration depth, the

curves begin to fall and asymptotically approach the hardness or Young’s modulus of

the Si substrate. The indentation depth is directly related to the magnitude of the

indenter load and the relationship between the load propagation, dP/dh and the

combined Young’s modulus for a unload cycle is defined as

απ

tanh2E2dhdP

t*=

(33)

where dP/dh is the load propagation and tan α is the slope of the unloading curve, E*

is the combined Young’s modulus and ht is the penetration depth.

The load propagation as a function of penetration depth for a-C:H samples

deposited at different bias is shown in Fig. 46, which shows dP/dh at the initial point

of contact for samples deposited at high bias (-400 V) is higher than for samples

deposited at low bias (-250 V).

Fig. 46. Load propagation dP/dh vs. penetration depth h, for a-C:H samples

deposited at different bias.

78

At the depths of ~6 µm and over, the dP/dh is equal for all examined samples as at

that depth the Si substrate is being probed and contributes as much to the elastic-

plastic response as the film. The summary of H, E and H/E ratio results for all a-C:H

films studied are shown in Table 5. The H/E ratio increased by a small amount for

films deposited under bias above -350 V. This ratio H/E is 0.1 for most materials and

is the ratio for natural diamond 1.

Table 5

Nanoindentation results for a-C:H samples produced under various substrate bias.

The H/E ratio accounts for relative changes in hardness that are related to density,

whereas the magnitude of E relates to interatomic bonding forces325 and is

proportional to the slope of the interatomic force-separation curve at the equilibrium

spacing

0rdrdFE ⎟

⎠⎞

⎜⎝⎛∝

(34)

Where, dF/dr is slope of the load-displacement curve and r0 is the equilibrium

spacing. Therefore, the ratio H/E characterises the physical response of an atomic

lattice to an applied external force. The ratio H/E relates to the fracture strength

which is found to be slightly higher for a-C:H samples deposited at higher bias. This

small, but based on the data above, significant difference indicates that inter-domain

stress is elevated in films deposited at higher bias and these films contain stress

Bias,

V

Hardness,

± 1 GPa

Young’s

modulus,

± 3 GPa

EH

ratio

±0.007

- 250 15.2 174 0.087

- 300 15.6 180 0.087

- 350 18.0 187 0.096

- 400 18.6 190 0.098

79

raisers. The values for hardness and Young’s modulus obtained in our measurements

were found to be in good agreement with results reported in the literature 2,326,327.

The next sections look at the bonding in these a-C:H films deposited under

different bias conditions, in an attempt to understand whether differences in the sp3

content are responsible for the differences in mechanical behaviour.

5.2 X-ray C1s results and discussion for a-C:H

The analysis of the XPS spectra is performed by decomposition of the C1s core

binding energy (BE) spectra into two constituent functions corresponding to sp2 and

sp3 carbon hybridised states as described in Section 4.3. The comparative BE spectral

functions of graphite and natural diamond234,247 were used to calculate the sp3/sp2

ratio based on their respective area ratios233,248. The raw data curves before

background subtraction for all the samples analysed in this Chapter are shown in Fig.

47. For a-C:H films examined the main core-level C1s peak shifted by a small

fraction of 0.07 eV to higher BE from 284.58 eV to 284.65 eV for a-C:H samples

fabricated at higher bias (-400 V) as shown in Fig. 47.

280 282 284 286 288 290 292Binding energy, eV

Rel

ativ

e in

tens

ity (a

rb. u

nits

) 400 V

350 V

250 V

200 V

284.65 eV

284.58 eV

C 1s 400 V

350 V

300 V

250 V

280 282 284 286 288 290 292Binding energy, eV

Rel

ativ

e in

tens

ity (a

rb. u

nits

) 400 V

350 V

250 V

200 V

284.65 eV

284.58 eV

C 1s 400 V

350 V

300 V

250 V

Fig. 47. The C1s spectra of a-C:H films fabricated under varying substrate bias.

80

Surface charging was not observed during the data acquisition. Due to the

environmental oxidation two additional peaks were introduced into the fitting of the

main C1s spectra, a C-O bond at 287 eV and a C=O bond at 288.4 eV30,234. As

discussed in Chapter 4.3 and 4.4.1 above, the P VII function was used for fitting the

sp2 and sp3 components with the P VII parameter M of 2 for sp3 and 3 for sp2 phase.

Lorentzian line-shapes were used for C-O and C=O peaks. As further discussed in

Chapter 4.3, a constraint on the binding energy difference between the sp2 and sp3

peaks of 0.8 < ΔBE < 0.9 eV was used during the fitting procedure. The XPS C1s

spectrum for an a-C:H film is shown in Fig. 48; the parameters of the fit for each

sample are given in Table 6. The fitted XPS C1s spectra for all a-C:H films, i.e.

fabricated under -250, -300 and at -350 V bias are presented in Appendix 2. The

sp3/sp2 ratio was found to be 0.28 ± 0.02 for low bias (-250 V) and 0.28 ± 0.02 for

high bias (-400 V) samples indicating that any change in sp3 fraction with the

substrate bias is undetectable.

282 283 284 285 286 287 288 289 290 291 292Binding energy, eV

Rel

ativ

e in

tens

ity (a

rb. u

nits

)

sp 3sp 2

C-O C=O

Fig. 48. The fitted XPS C1s spectrum of an a-C:H film fabricated under – 400 V.

The sp2 peak BE shifted from 284.63 to 284.59 eV for higher substrate bias, whereas

the BE of the sp3 fraction was unchanged and within the range 285.47 ± 0.02 eV for -

250 V bias to 285.50 ± 0.02 eV for -400 V bias. ∆BE changed from 0.86 to 0.91 eV,

which is not significant given the fitting uncertainties.

81

Table 6

The detailed XPS C1s results for samples produced under various substrate bias. The

fitting uncertainties for hybridised states are ± 0.02 eV, for the sp3/sp2 ± 0.018 and

for ∆BE is ± 0.03 eV.

*this value is calculated from the curve fit, but the uncertainty, based on the variation

of the peak positions is ± 0.02. Therefore all values quoted above are 0.28 ± 0.02.

The increase of ∆BE for samples fabricated at higher bias indicates that the

FWHMsp3 is somewhat larger than FWHMsp2 and the structural disorder broadens the

core level shift 233,234,247,248. The ∆BE was found to be within a range presented in the

recent publications233,234,239,328,329.

5.3 Raman spectroscopy results and discussion for a-C:H

Fig. 49 shows MW Raman spectra obtained using 633, 532, 325 and 244 nm

excitation wavelength for a a-C:H film sample fabricated under -400 V bias. All

spectra were fitted with the common G and D peaks (visible excitation), the T peak

The P VII line-shape was used with P VII width M of 3 for the D and D* peaks and

at the value of 5 for the G peak; the uncertainty of the peak position fitting was ±

0.75 cm-1. Fig. 49 shows the dispersion of all fitted bands at higher photon excitation

energy due to phonon confinement formalism259,273,280. The analytical results for a-

Bias, V C1s, eV sp2, eV sp3, eV sp3/sp2 * ∆BE

- 250 284.58 284.63 285.47 0.281 0.86

- 300 284.60 284.62 285.47 0.282 0.87

- 350 284.65 284.62 285.49 0.284 0.87

- 400 284.65 284.59 285.50 0.285 0.91

82

C:H samples under selected Raman excitation wavelength are presented in the

following Sections.

The 633 nm Raman results

The results showed a broad G peak centred at approximately 1520 cm-1. Regardless

of negative bias voltage, the position of the G band remained unchanged at 1520 cm-1

for all films studied. The FWHMG increased from 195 to 197 cm-1 for films deposited

at higher bias (-400 V). The D peak was positioned at 1350 cm-1 and the position was

not changed at any bias voltage. The FWHMD increased from 210 to 220 cm-1 for

films deposited at higher bias indicating a size reduction for the sp2 aromatic

domains. Finding the D* peak was unexpected since shoulder peaks are rarely

observed in Raman spectra, however a peak located at approximately 1150 cm-1 has

been used with the two main G and D peaks in the analysis is spectra by other

workers330-333. Peaks at a similar positions were found in poorly organised

amorphous carbons and Ferrari331 has suggested that a peak at this frequency might

be a trans-polyacetylene (TPA) inclusions, whereas Yan et al332 attributed this peak

to unbound hydrogen in a-C:H. We found the D* peak centred at 1193 cm-1 and the

peak position did not change at any bias voltage. FWHMD* remained unchanged at

180 - 181 cm-1. The ratio of the peak intensities for the D and G peaks, I(D)/I(G),

decreased from 0.66 to 0.62 for films deposited at higher bias showing that samples

fabricated at lower bias contain fewer sp2 aromatic rings and these are of smaller

size. The ratio I(D*)/(I(D) + I(G)) increased form 0.15 to 0.17 for high bias films.

Overall, the analysis of 633 nm Raman indicated a weak reduction of aromatisity due

to the increased bias.

The 532 nm Raman results

The G peak position moved from 1537 cm-1 to 1532 cm-1 for films produced at high

bias indicating the increase of sp2 olefinic sites. The FWHMG decreased from 189

cm-1 to 193 cm-1 for films deposited at higher bias. The D peak was positioned at

approximately 1378 – 1380 cm-1 and the FWHMD changed from 228 cm-1 to

232 cm-1 indicating that the size of sp2 aromatic rings decreased in higher bias films.

The ratio I(D)/I(G) decreased from 0.48 to 0.43 for films film deposited at higher

bias, as in the 633 nm Raman measurements, confirming the decrease of aromaticity

at high substrate bias. The D* peak at 1198 – 1199 cm-1 and ratio I(D*)/(I(G) + I(D))

83

changed from 0.13 to 0.12. FWHMD* was largely unchanged at 146 cm-1. The 532

nm Raman analysis confirmed the conclusions for the sp2 phase changes as derived

from the 633 nm Raman measurements.

Fig. 49. MW Raman spectrum of an a-C:H film fabricated under -400 V bias.

84

The 325 nm Raman

Again the G peak position moved from 1584 cm-1 to 1589 cm-1 while the FWHMG

increased from 122 cm-1 to 126 cm-1 indicating a higher degree of disorder of the sp2

phase. The D peak is present in UV indicating that there is some amount of the sp2

fraction arranged in aromatic rings. The D peak position was at constant 1454 cm-1.

for all samples. The sp3 sites were visible as indicated by a small T peak at

approximately 1031 cm-1. There was a negligible increase of the ratio I(T)/I(G) from

0.040 to 0.041 for films deposited at high bias. This ratio decreased from 0.27 to 0.24

for films fabricated at high bias indicating the fall in sp2 aromaticity. The D* peak

position remained at approximately 1200 - 1201 cm-1, a frequency similar to that

found in the 532 nm Raman measurements. The contributions from the D* peak were

very minor in however in order to obtain a quality fitting its contributions had to be

included. The ratio I(D*)/(I(G) + I(D)) was approximately 0.02 for films deposited at

high and low bias.

The 244 nm Raman

The 244 nm Raman is ideally suited for evaluation of both the sp3 and sp2 phases in

DLC271,274,334. The G peak position moved from 1589 cm-1 to 1595 cm-1 with the

increase of bias and the FWHMG increased to 131 cm-1 from 126 cm-1 for films

fabricated at higher bias. The D peak was positioned at 1460 cm-1 and the FWHMD

increased by 5 cm-1 for higher substrate bias films. The T peak was positioned at

1015 cm-1, and the FWHMT was 110 cm-1 for films deposited at higher bias. The ratio

I(T)/I(G) was found to be 0.038 for all films confirming that amount of the sp3 phase

was equal between the samples examined. The D* peak position was at 1205 - 1206

cm-1. The ratio I(D*)/(I(G) + I(D)) was 0.014 for all examined films. The D* peak

intensity strongly decreased at this excitation relative to the 325 nm. It was not

possible to quantitatively estimate the amount of the sp3 fraction for a-C:H films

using the 244 nm results as, for example, it is possible for hydrogen free DLC, since

the T peak is only sensitive to C-C sp3 bonding and not to C-H sp3 bonding259,260,262.

High intensity and definition of the T peak at both the 325 and 244 nm exitations

indicate that there is a large fraction of C-C sp3 bonded sites in the a-C:H films

studied.

85

Overall, MW Raman analysis indicated that all of the a-C:H films examined

display essentially the same amount of C-C sp3 phase indicated by the intensity and

position of the T peak in the UV excitation Raman spectra. The arrangement of the

sp2 phase in all a-C:H samples was found to be different. The lower substrate bias (-

250 V) a-C:H samples displayed more sp2 aromatic clusters and less sp2 olefinic sites

compared to the higher biased (-400 V) samples. This is revealed by the D and G

peak trends. The sp2 olefinic sites in higher biased samples are arranged in smaller

domains, deduced from the relative shifts of the G peak position and changes of the

FWHMG.

The D* peak39 is certainly of the sp2 origin335 as its intensity is greatly reduced

at higher laser excitation energies. In a separate work336 we found that the D* peak

has a companion mode at approximately 1450 cm-1 and the 1450 cm-1 band becomes

more visible at low excitation energy. Indeed the origin of the D* peak is mot likely

due to TPA like clusters331. The TPA inclusions in nanocrystalline diamond were

first believed to be of the sp3 origin256,337,338, however recently Ferrari and

Robertson331 proved the peak assignment to TPA and confirmed this is of the sp2

phase.

The Raman in the visible range (532 nm) was found to be better suited for the

sp2 phase analysis as compared to the 633 nm wavelength by giving more defined

shapes for the core D and G bands and their Raman shifts.

5.3.1 Rayleigh scattering results for a-C:H

The 532 nm Raman line was used for the measurement of Rayleigh

scattering (intensity, IBS) discussed in Section 4.4.2. IBS was extracted from the broad

asymmetric peak centred at 0 cm-1 which was deconvoluted into three constituent

Lorentzian peaks at -1.25, 0 and 1.25 cm-1. The intensity of the middle peak centred

at 0 cm-1 was used to determine the final Rayleigh scattering intensity, IBS. Fig. 50

86

shows the intensity of the Rayleigh scattering line expressed in arbitrary units

relative to the substrate bias of fabricated a-C:H films. It illustrates that the IBS almost

linearly relates to the increase in substrate bias used during deposition of a-C:H

films. The films fabricated at higher bias are of superior mechanical properties

(Section 5.1) and as seen in Section 5.1 above, where the H/E ratio shows these films

are of higher density, and the hardness of the films deposited at -350V and -400V

substrate bias is significantly higher than the films deposited with lower substrate

bias. Therefore, the extracted value of IBS can also be used to monitor the relative

density changes in DLC films.

200

250

300

350

400

450

500

225 250 275 300 325 350 375 400 425Negative bias, V

Ray

leig

h in

tens

ity (a

rb. u

nits

)

Fig. 50. Relationship between the height of 532 nm scattered Rayleigh line for a-C:H

samples fabricated under different bias.

While it appears easy to obtain the backscattering results and to interpret them, care

needs to be taken when drawing conclusions from the Rayleigh measurements since

the IBS is known to relate to the surface roughness50,211,339-342 of the sample and is also

directly proportional to the distance from the focal point of the laser source to the

sample surface. The surface roughness of all examined a-C:H films was found to be

approximately the same as evidenced by examining Fig. 57 – 60 of Section 5.6. The

roughness was found to be approximately 10 - 20 nm.

87

5.4 IR spectroscopy results for a-C:H

IR absorption spectra were obtained for all fabricated a-C:H films. The

spectrum of an a-C:H sample in 650 – 4000 cm-1 region is shown in Fig. 51. The IR

spectra for all a-C:H films fabricated under different bias voltage is displayed in

Appendix 3. All IR spectra were analysed in 2700 – 3200 cm-1 region (the region

where C-H stretching vibrations are evident) and in 1000 – 1700 cm-1 region

(looking at C-C and C-H bending and rotational modes)2,291,292,343. The assignments

of C-H stretching modes are summarised in Table 7.

60010001400180022002600300034003800Frequency, cm-1

Abs

orpt

ion

(arb

. uni

ts)

Fig. 51. IR absorption spectrum of a-C:H sample deposited at -250 V.

The actual IR absorption spectra in 2700 – 3200 cm-1 region for a-C:H samples

fabricated at -250 V and -400 V is shown in Fig. 52. A contribution from each

individual C-H stretching group was evaluated based on a group’s respective square

area, A. A was calculated by taking the intensity over the FWHM for the selected

peak. Prior to fitting, a linear background was subtracted. The uncertainty of the peak

position were ± 0.25 cm-1. The results of the spectral fittings for 2700 – 3200 cm-1

region are summarised in Table 8. These The results indicate the strong increase in

spectral contributions from the sp2 group with vibrational frequencies of 3130 and

2995 cm-1 at higher bias due to the bond saturation. At the same time the

contributions from the sp3 peaks at 2920 and 2855 cm-1 have increased. The

intensities and FWHMs of the sp2 peaks positioned at 3085 and 3035 cm-1 decreased,

88

however the contributions from CH2 peak at 3085 cm-1 have decreased to a

negligible amount of 0.7 % at high bias voltage.

Table 7

Assignments of a-C:H IR vibrational frequencies in the 2700 – 3200 cm-1 region; C-

H stretching vibrations 344-346.

Frequency,

cm-1

Assignment

State Group Arrangement Symmetry

3130 sp2 C=C-H unsaturated aromatic Asymmetric

3085 sp2 CH2 unsaturated olefinic Asymmetric

3035 sp2 =C-H saturated aromatic Asymmetric

2995 sp2 =C-H saturated olefinic Symmetric

2975 sp3 -CH3 saturated olefinic Asymmetric

2920 sp3 =C-H and =CH2 saturated olefinic Asymmetric

2855 sp3 =CH2 saturated olefinic Symmetric

The –CH3 sp3 peak at 2975 cm-1 decreased strongly at higher bias. The observed

relative increase in saturation of the sp2 bonded groups and, in particular, the =C-H

symmetrical group, at higher bias was attributed to the higher energy of incident C

ions and the concurrent increase of the etching rate by hydrogen ions. This process

also related to the changes observed for the sp3 =CH2 group. It is known for a-C:H

films to contain all hydrogenated sp3 bonded groups, however some sp2 groups are

often found to be un-hydrogenated as reported by Tamor et al180 and Donnet et al347.

Owing to that observation we refrained from the estimating the relative sp3 content

from the IR results. The UV Raman analysis (Section 5.3) showed the defined T peak

of high intensity confirming that indeed C-C sp3 bonds are present in a large amount.

Additional information about the morphology of a-C:H samples was obtained

from N- IR absorption spectra in 1000 – 1700 cm-1 region corresponding to C-H in-

plane bending and some C=C stretching vibrations. The assignment of a-C:H active

frequencies in 1000 – 1700 cm-1 excitation region is shown in Table 9.

89

Fig. 52. The deconvoluted IR absorption spectra in 2700 – 3200 cm-1 region for a-

C:H samples deposited at -250 and -400 V

Table 8

Calculated relative peak areas, A as a % of total peak area for C-H stretching groups

in 2700 – 3200 cm-1 region.

Sample 3130 cm-1

sp2

3085 cm-1

sp2

3035 cm-1

sp2

2995 cm-1

sp2

2975 cm-1

sp3

2920 cm-1

sp3

2855 cm-1

sp3

-250 V 2.8 9.9 10.7 7.8 6.9 38 24

-300 V 3.3 5.5 9.6 8.8 5.7 42 25

-350 V 4.3 4.1 5.8 9.3 4.6 46 26

-400 V 4.5 0.7 4.9 10.8 3.0 49 27

90

Table 9

The N-IR vibrational frequencies in 1050 – 1700 cm-1 region for a-C:H 343.

Frequency, cm-1 Assignment

Origin Arrangement/ Symmetry Mode

1600 sp2 -C=C- olefinic, asym. stretch

1470 - 1420 sp3 -CH3 olefinic, asym./sym. in-plane bend

1450 sp3 =CH2 olefinic, sym. in-plane bend

1310 - 1290 sp2 =C-H olefinic, asym. in-plane bend

~ 1100 sp2 =C-O-C= aromatic, asym. in-plane bend

The N-IR spectra were analysed based on contributions from constituent peaks, as

for the IR spectra above. The detailed analysis of all constituent groups was not

possible due to presence of large unresolved absorption bands at approximately at

1400 – 1550 cm-1 and above the 1650 cm-1 region. Fig. 53 shows that the intensity of

the 1600 cm-1 peak corresponding to the C=C sp2 stretching mode increased at

higher bias indicating the increase of sp2 phase disorder (de-aromatisation).

Fig. 53. Comparative IR absorption spectra in 1050 – 1700 cm-1 region for a-C:H

samples fabricated at -250 V and -400 V substrate bias.

91

Peaks positioned in a low resolution region of 1400 – 1550 cm-1 belong to the sp3 -

CH3 and =CH2 groups. The intensity of a well-defined =C-H sp2 peak centred at

around 1300 cm-1 increased with increasing bias. There is also a C-O contamination

band at 1100 cm-1 343, but this is not discussed since it is irrelevant. The analysis of

N-IR region complimented the information that was obtained from analysing the IR

results.

5.5 Band gap measurements for a-C:H. Results and

discussion of Tauc gap vs. surface conduction gap.

Fig. 54 shows extrapolation of fundamental absorption curves in N-IR range

for all examined a-C:H samples from which the Tauc gap values were measured

(Section 4.6).

0

10

20

30

40

50

60

70

80

90

100

940 965 990 1015 1040 1065 1090 1115 1140 1165 1190

Wavelength, nm

Abs

orbt

ion,

%

-400 V-350 V-300 V-250 V

Fig. 54. Extrapolation of N-IR absorption spectra for a-C:H samples.

The absorption wavelength for a-C:H samples fabricated at -400, -350, -300 and -250

V were 1024, 999, 988 and 969 cm-1 respectively and corresponded to band gaps of

1.2, 1.24, 1.25 and 1.28 eV respectively. Fig. 55 shows the Tauc gap, Eg and the

92

surface conduction gap, ESC measured by STS (Section 4.6, 4.6.1) as a function of

the suibstrate bias during deposition.

1.15

1.2

1.25

1.3

1.35

1.4

200 250 300 350 400 450Negative bias, V

Ban

d ga

p, e

V E SC

E g

Fig. 55. The variation of the Tauc, Eg and surface conduction gap, ESC with bias

voltage for examined a-C:H films.

Clearly the Eg does not correlate well with ESC, and indeed there appears to be an

anti-correlation – as Eg increases, Esc decreases. This anti-correlation can be easily

explained taking in to the account the nature of the respective probing techniques

used. The Tauc gap uses IR radiation and probes the bulk of the DLC sample,

whereas the other (the STS), is surface sensitive. It is well known that the band gap

in DLC is controlled by the π – π* gap of the sp2 hybridized phase and the reduction

or of the sp2 fraction causes the band gap to become wider (Section 4.6)301,348,349. Not

only the actual amount of the sp2 fraction, but also the ordering (bond angle disorder,

small domains) of the sp2 fraction is also known to affect the band gap301,348,349. It

was determined from the Raman and IR results that the sp2 sites in a-C:H films

fabricated at higher bias are primarily olefinic with high degree of bonding disorder.

The band gap of sp2 domains is controlled by perturbation of the π states and inherent

tail bands lead to narrowing of the conduction gap. This is illustrated in Fig. 56. Fig.

56 shows that the π - π* conduction gap can become narrower relative to the

unchanging σ - σ* gap. For that to occur there is no need for the π states to become

reduced, the π states can become perturbed by low energy of incoming C atoms

during the DLC formation.

93

Fig. 56. A) Schematics VDOS of DLC and B) Perturbation of π states is shown.

Lower Tauc gap values in higher biased a-C:H samples indicate that the arrangement

of sp2 phase changed independently to a minor degree relative to the sp3 fraction. The

ESC gap however, increased with the increase of negative bias indicating the raised

conductivity of the sp2 rich uppermost layer. The uppermost layer of a growing film

is not affected by the change in the VDOS of the bulk but only at the surface.

Therefore it only reflects VDOS changes corresponding to the changing energy of

incoming ions. The energy of C ions is elevated when bias is increased resulting the

top layer to become more of the sp2.

5.6 SEM images for a-C:H

All a-C:H films (frontal and lateral surfaces) were examined by SEM and the

films were found to be smooth with the surface roughness of less than 20 nm, see

Fig. 57 - 60. Fig. 57 shows the surface of an a-C:H film fabricated under -250 V bias

voltage, Fig. 58 displays the surface of an a-C:H film produced under -300 V bias,

Fig. 59 shows a -300 V and Fig. 60 a film fabricated under -400 V.

94

Fig. 57. The surface of an a-C:H film fabricated under -250 V bias.

Fig. 58. The surface of an a-C:H film fabricated under -300 V bias.

95

Fig. 59. The surface of an a-C:H film fabricated under -350 V bias.

Fig. 60. The surface of an a-C:H film fabricated under -400 V bias.

96

That value was later confirmed by atomic force microscopy. Film thickness of a-C:H

samples on Si was found to be 9 ± 0.02 µm. Fig. 61 shows the lateral image of an a-

C:H film fabricated at -300 V negative bias.

Fig. 61. Lateral image of an a-C:H film on Si substrate.

5.7 Discussion of obtained results for a-C:H

The fabricated a-C:H films were found to be of a narrow range mechanical

properties350,351 and films fabricated at higher substrate bias (-350V, -400 V) display

superior mechanical properties as compared to the films fabricated at lower substrate

bias (-250 V, -300 V). The actual differences are very small. For film hardness the

difference is an increase approximately 3.4 GPa (20%) and for Young’s modulus an

increase of 16 GPa (10%) over films deposited at low substrate bias, however, the

97

differences can be considered significant with the uncertainties in hardness being ± 1

GPa and in Young’s modulus ± 3 GPa. The nanoindentation measurements acquiring

the mechanical properties results were performed using current well established

methodology (over 10000 load-unload indentation curves for each sample). The

small differences in mechanical properties between high and low bias films are also

reflected in comparable sp3 content for either of these films found to be

approximately 28 %. Films fabricated under higher bias are of higher structural

disorder as indicated by wider ∆BE shift39,350 observed in the XPS measurements. By

structural disorder we mean bond length and bond angle disorder and arrangement of

carbon sites into olefinic-like groups over aromatic-like. This structural disorder in a-

C:H films when studied in detail by IR spectroscopy351 shows that indeed the

increase of bias contributes to the process of hydrogen abstraction (de-

hydrogenation) for both sp2 and sp3 constituents. This is primarily evidenced by

decreased contributions from the tetrahedral –CH3 sp3 group together with other sp3

and sp2 bonding groups at higher bias. The increase of bias was found beneficial for

formation of secondary, tetrahedral =C-H and =CH2 sp3 groups, however in this case

the changes can be attributed to the increased energy of H and C ions through the

physical bombardment process. This physical bombardment process also contributed

to inhomogeneity in fabricated films. The inhomogeneity was found to bias

dependent as identified by using MW Raman, whereas the STS and N-IR

measurements showed the uppermost layer in fabricated films was more sp2 like with

high perturbation of the sp2 π states. The analysis of MW Raman results confirmed

that sp2 sites at higher bias become more olefinic with shorter chain lengths and

smaller domains, while only minor changes in sp3 sites occur due to the bias

variation.

Overall, all of the analytical techniques used demonstrate that at a given

condition where sp3 constituent is equal among all a-C:H samples produced, it is

ordering of sp2 component that finally determine the mechanical properties of the

films and, to the extent their electronic properties.

98

99

Chapter 6

6. Characterisation of hydrogen free DLC

fabricated using the RIBSD

Fabrication of DLC films using the RIBSD technique is critically examined in

this Chapter concentrating on capabilities and limitations of the technique for

producing quality (medium-high sp3 content) DLC samples. The microstructure of

fabricated hydrogen free a-C and DLC films was examined with a focus on

determining only the sp3 component, for that purpose UV (325 nm) Raman and XPS

C1s analytical techniques were used.

6.1 UV Raman analysis of hydrogen free DLC

Films that were obtained as a result of the RIBSD experiments were

categorised into three groups:

• The first are the undeveloped carbon films on Si substrate, the SiC films,

these normally provide an intermediate, transitional layer for further

growth of nucleating carbon film when deposited on Si. Such SiC films are

very thin, with thickness of less than 15 nm352-355.

• The second are the a-C films with low sp3 content (less than 7%), these

films are produced when the energy of C ions is low. The a-C films are

graphite-like with high degree of aromatic ordering. The a-C are known to

be devoid of characteristic spectral features259,262,356 of DLC and when

examined by UV Raman these films show no T peak. The a-C films are

now considered unsuitable for applications in optics, electronics or

tribology.

100

• The third films group produced are the actual DLC films with substantial

amount of the sp3 content as also evidenced by the strong T peak in UV

Raman spectra. Here we need to make an important remark that DLC

produced using the RIBSD are the films containing certain amounts of the

sp3 inclusions. However, these DLC are very inhomogeneous owing to the

nature of the target sputtering process and most certainly include graphitic

clusters, nano-sized graphite particles and contain traces of sputter atoms

(Ar or Xe) that are known to become impregnated during such deposition

methods125,357,358.

The results obtained using the RIBSD are summarised in Table 10 that shows

SiC, a-C and DLC films fabricated using varying ion energies, ions (Ar and Xe) and

the sputtering geometry. The data entered in the Table 10 is the result of analysis of

all UV Raman spectra (see Appendix 4) for the RIBSD fabricated samples. The

DLC, SiC and a-C films were categorised into three group based on their inherent

UV Raman spectra as discussed above.

Table 10

Films fabricated using the RIBSD at varying Ar and Xe ion beam energies and

target/substrate sputtering geometry. “--" shows the experiments were no performed.

Ion

energy

keV

Target and substrate to the ion beam axis sputtering angles, αt : αs (°)

15° : 0°

15° : 10°

30° : 0°

30° : 10°

45° : 0°

45° : 10°

0.2 --- --- SiC --- --- ---

0.4 --- SiC a-C SiC --- ---

0.6 a-C SiC DLC a-C a-C SiC

0.8 DLC --- DLC --- DLC ---

1.0 --- a-C DLC a-C --- a-C

1.2 --- --- DLC a-C --- ---

101

The RIBSD results indicate that in cases where the substrate was positioned at

grazing angle, αs, of 10°, there were no DLC films formed, but only SiC and a-C.

Fig. 62 shows UV Raman spectra for two a-C films produced using Ar and Xe ions

with ion energy of 1.2 keV bombarding the target at 30° and the substrate at 10° (αt, :

αs is 30° : 10°).

Fig. 62. UV Raman spectra of a-C films fabricated at sputtering angles of αt 30° and

αs 10° and sputtering ion energy of 1.2 keV for Ar and Xe ions.

900 1000 1100 1200 1300 1400 1500 1600 1700 1800

Wavenumber, cm-1

Inte

nsity

(arb

. uni

ts)

T

D

G

Fig. 63. Fitted UV Raman spectra of an DLC film fabricated using 1.0 keV Xe ions.

The αt, : αs was 30° : 0°.

102

Even without the detailed peak fitting it is evident that these a-C films are sp2

polymer-like with large aromatic content as evidenced by the large D band. No T

band260 indicates that these a-C films are devoid of the sp3 content. Contributions

from the Si substrate are seen in 930 – 980 cm-1 range. This indicates that the

deposited a-C films are very thin, perhaps of few tens of nanometers thick and

porous. The UV Raman spectrum for the RIBSD fabricated DLC film is shown in

Fig. 63. Fig. 63 shows that the T band is clearly visible and the spectrum overall is

dominated by the G peak indicating that DLC films are sp2 olefinic with the sp3

imbedded sites. The I(D)/I(G) and the I(T)/I(G) ratios were determined from the UV

Raman spectra for all films and were referenced against the ion beam energy, keV as

shown in Fig. 64, and Fig. 65.

0

0.2

0.4

0.6

0.8

1

1.2

1.4

0 0.2 0.4 0.6 0.8 1I(D)/I(G)

Ion

ener

gy, k

eV

30_0

15_0

15_10

30_10

45_0

45_10

Fig. 64. The relationship between the ion bombardment energy and I(D)/I(G) ratio

for the RIBSD fabricated films. The legend shows the angles of target and substrate

as αt_ αs

The analysed Raman spectra showed almost no distinction between the Ar and

Xe fabricated films, with the exception of varying film thickness (films were thicker

when sputtered by Xe), therefore data for both sputter ions was presented together in

these figures. In Fig. 64 and 65, the legend gives the substrate and target angles wrt

to the incident ion beam as αt_αs. All UV Raman spectra were fitted with the P VII

line and the P VII M coefficient was 5 for the G peak and 3 for the D and the T

103

peaks. Fig. 65 shows that the increase of ion bombardment energy for αt, : αs

combination of 30° : 0° reduces the number of sp2 aromatic sites as indicated by the

decrease of the I(D)/I(G) ratio.

0

0.2

0.4

0.6

0.8

1

1.2

1.4

0 0.05 0.1 0.15 0.2 0.25I(T)/I(G)

Ion

ener

gy, k

eV30_0

30_10

15_0

15_10

45_0

45_10

Fig. 65. The relationship between the ion bombardment energy and I(T)/I(G) ratio for

the RIBSD fabricated films.

The value of the I(T)/I(G) ratio can be used as a guidance to distinguish

between DLC films of low sp3 content from the films with the high content using the

empirical data for I(T)/I(G) ratio and the T position, cm-1 derived by Ferrari and

Robertson260 specifically for UV Raman measurements. That is the I(T)/I(G) < 0.2

corresponds to sp3 content < 30 %260. The I(T)/I(G) ratio for majority of the studied

DLC samples was found to be in 0.13 – 0.20 range, see Fig. 65. However, these films

at a given I(T)/I(G) ratio also show the T band position that falls within 1065 – 1070

cm-1 range. The T position corresponds to sp3 content of up to approximately 50

%259,260.

The application of the impinging bombardment, that is when the substrate

surface is exposed to the ion directed ion beam flux, at αt, : αs combination of 15° :

10°, 30° : 10° and 45° : 10° were found to be unfavourable for formation of DLC and

at these sputtering combinations the I(D)/I(G) ratio was also quite high. Correlating

the informaiton in Fig. 64 and Fig. 65 affirms that the αt, : αs combination of 30° : 0°

104

is better suited for DLC fabrication as evidenced by the strong T peak as compared to

configurations when αs>0°. Apart from the configuration αt=15°, αs=0°, that is

decreasing the target to the ion beam incident angle, no T peak presence was found

for other films studied.

6.2 C1s X-ray Photoelectron Spectroscopy analysis for

hydrogen free DLC

The analysis of C1s XPS spectra for a-C and DLC films was performed using

the approach233,248 described in Sections 4.3 and 5.2. During XPS data acquisition

strong surface charge accumulation for sp2 rich a-C films was observed and there

was no charge accumulation for DLC samples. The differences in the obtained C1s

spectra between the sp2 rich a-C films and the DLC films are illustrated in Fig. 66 A

(a-C sp2 rich) and B (DLC). Fig. 66 A shows that the BE position for a-C film is also

affected by surface charging. The XPS spectra were fitted with the P VII line

functions with coefficient M of 3 for the sp2 and the sp3 components and Lorentzian

were used for the C-O and C=O peaks232,247. Peak positions for oxides were 286.9

eV (C-O) and 289.1 eV (C=O) for an a-C and 286.43 eV (C-O) and 288.1 eV (C=O)

for a DLC sample. The sp2 component was found positioned at BE of 285.25 eV with

FWHMsp2 = 1.26 eV for the a-C film and at 284.15 eV with FWHMsp2 = 1.13 eV for

the DLC. For the sp3 constituent the BE was 285.01 eV with FWHMsp3 = 1.45 eV for

the a-C film and BE of 286.15 eV with FWHMsp3 =1.56 eV for the DLC. The

uncertainly of the fitting was ± 0.02 eV. The ∆BE233,234,329 for the best fit were 0.86

eV for the a-C and 0.9 eV for the DLC, within the theoretical constraint discussed in

Chapter 4.2. The sp3/sp2 ratio was determined to be 0.06 (6 % sp3) for the a-C sample

(which angles) and 0.37 (37 % sp3) for the DLC film (angles).

105

Fig. 66 A: sp2 rich a-C film fabricated using Ar bombardment at 1.0 keV and αt, : αs

of 45° : 10°, B a DLC film fabricated using Xe ions at 1.0 keV and αt, : αs of 30° : 0°.

The summary of sp3 content expressed in %, obtained for all the a-C and DLC films

fabricated using sputtering with Ar and Xe ions is shown in Table 11. In was noted

earlier that the UV Raman spectra analysis for these films showed almost no

distinction between the Ar and Xe fabricated films, apart from the varying film

thickness (Section 6.1). The results shown in Table 11 also indicate that indeed the

variation of the sp2 component for Ar or Xe fabricated films varies insignificantly.

The sp3 amount found in Xe sputtered DLC films is 45%, while for Ar sputtered films

this value is slightly lower at 43%. All sp3 rich DLC films were deposited when the

αt, : αs combination were 30° : 0°, 45° : 0° and 15° : 0°, that is for different ion beam

to target incidence angles (αt), but with the substrate positioned parallel to the axis of

the incoming ions. When grazing bombardment was used (incident beam to

subbstrate angle of 10°) no DLC films were formed.

When the estimate for sp3 content obtained using XPS C1s measurements was

correlated with the T peak intensity, I(T) of UV Raman measurements the following

relationship is obtained, Fig. 67.

106

Table 11

The sp3 content, ± 1.5 % of a-C and DLC films fabricated using the RIBSD method

with Ar and Xe ions.

Ion

energy

keV

Target and substrate to the ion beam axis sputtering angles, αt : αs

15° : 0°

15° : 10°

30° : 0°

30° : 10°

45° : 0°

45° : 10°

0.2 --- --- SiC --- --- ---

0.4 --- SiC ≤ 5 (Ar & Xe)

SiC --- ---

0.6 ≤ 7

(Ar & Xe) SiC 16 (Ar)

20 (Xe)

≤ 7

(Ar & Xe) ≤ 8

(Ar & Xe) SiC

0.8 18 (Ar)

21 (Xe) --- 32 (Ar)

34 (Xe) ---

22 (Ar)

24 (Xe) ---

1.0 --- ≤ 7

(Ar & Xe) 39 (Ar) 41 (Xe)

≤ 8 (Ar & Xe)

--- ≤ 7 (Ar & Xe)

1.2 --- --- 43 (Ar) 45 (Xe)

≤ 8 (Ar & Xe)

--- ---

Two conclusions can be made from examining the relationship shown in Fig.

67. First, there is a clear, almost linear, dependence of the I(T) on the sp3 content in

DLC films. That is just about expected since the same relationship exists for DLC

films when probed by deeper UV Raman lasers like 244 or 215 nm259,262,331,356,359.

The second conclusion relates to the XPS C1s measurements that are known to be

only surface sensitive. Here however, the clear correlation with the UV Raman

results indicates that the XPS measurements can also be effectively used to

determine the sp3 content of the bulk.

107

Fig. 67 The relationship between the sp3 content (XPS C1s) in the RIBSD films and

the T peak intensity (325 nm UV Raman).

The use of XPS in DLC analysis has always been debatable since it is

believed152,360,361 that the content of the uppermost layer (that is less than 20 Å)

which is accessible to XPS is usually heavily oxidised or largely graphitic. In the

case of the RIBSD samples we find that the produced DLC (DLC not the a-C films)

films are not oxidised for reasons unknown to us. Indeed the some part of the signal

comes from the upper surface, about 20%241, but the rest of the signal originates from

the bulk of the film and the use of the bulk probing technique like Raman, and the

clear relationship between the XPS and Raman results, verifies that XPS does probe

more than just the surface layers. However,

6.3 Discussion of the RIBSD technique for DLC fabrication

By examining the results obtained for the RIBSD technique it becomes

apparent that impinging bombardment is not beneficial for formation of the sp3

fraction in carbon films. This can be understood by examining a range of

displacement energies for SiC, sp2 sites and the sp3 sites of carbon. First let us to

examine the incident Ar or Xe ions promoting formation of SiC.

108

• Incident Ar/Xe ions promote formation of SiC when impingement

bombardment is present

SiC forms a layer that provides nucleation for carbon films when Si, as in our

case, is used as a substrate material. SiC is first formed when energetic C ions come

in contact with the substrate surface and it is estimated that thickness of this

interstitial layer in majority of cases when DLC is deposited in Si is less than 3

nm135,354,362-366. There are, in fact, four minimum recoil damage energies required to

create displacements in SiC depending on the projectile to target combinations 367.

These are: 41 eV (C in Si), 35 eV (Si in Si), 24 eV (Si in C) and 20 eV (C in C)367. It

is clearly evident that these displacement energies are significantly lower than the

energy delivered by Ar or Xe projectile ions to the C and Si atoms on the surface and

should therefore promote the formation of mainly SiC or to a lesser extent the sp2

rich a-C. The average sputtered C atom energy at αt=30° is only 30 eV at 1.2 keV for

Ar ions; the sputtered C atom energy for Xe is even lower. Therefore, we postulate

that a secondary re-sputtering process is occurring when the substrate is positioned at

grazing angles to the ion beam flux.

• Ar/Xe ions sputter form the carbon target at lower (αt =15°) or higher (αt

=45°) incident angle

The reduction of the angle αt to 15° leads to the reduction of the total number

of projectile ions that reach the surface of the target due to the greatly reduced ion

flux plane that is only a half (49 %335,368) when compared to the angle αt = 30°. The

reduced sp3 content as compared to the films fabricated at αt : αs of 30° : 0°

configuration at identical bombardment energies is a rather unexpected finding, since

the use of a lower αt angle should lead to production of ejected C ions of higher

energy and therefore, to promote sp3 formation. However, this can be understood by

examining the fact that, at reduced αt angle, the angle at which the sputter ions recoil

and the angle of ejected C ions are very low6, and comparable to the original ion

incidence angle. This will lead to a greater probability for a C ion to approach the

substrate surface at low incident angles thus reducing its chances of "sticking" onto

the surface. There are also far fewer C ions available to form a film since few are

leaving the target (see Fig. 30 and 31 in Section 3.2). All this contributes to

diminished sp3 content in DLC films fabricated at αt : αs of 15° : 0°.

109

The sp3 content is also reduced in DLC films produced at C αt of 45° (αs of 0°).

In this case the number of projectile ions reaching the target surface is increased by

over 30 %335,368 when compared to the angle αt = 30°. However, higher values of αt

correspond to lower energy (per ion) in the sputtered C flux and thus, is not

favourable for sp3 formation (see Fig. 30 and 31 in Section 3.2).

• Ar/Xe ions ablate the forming film

The re-sputtering process is also owned to the innate geometry of the RIBSD

system. The angle of approach at which Ar and Xe ions ablate the Si surface is low

and the plasma density is low, therefore creation of a confined plasma environment

between the target surface and the substrate is hardly achievable. That is, in a

simplified view, the projectile noble gas ions just sweep away the C ions atoms

ejected form the target.

• Type of projectile ions

The secondary re-sputtering process can not be attributed only to the energy of

the incoming ions or the sputtering geometry but also to the type of projectile ions

used. That is, if not the noble gas (which are non reactive) ions were used but the C

ions at the same projectile angles, the probability of creating the sp3 bonding on Si

will be higher. If noble gas ions were bombarding the Si surface at normal angle and

C ions were also present within the vicinity of the surface, the probability of

achieving high sp3 content in DLC will be also much higher.

110

111

Chapter 7

7. Summary of the results and contributions to the

existing field of knowledge

The Chapter presents the summary for experimental and analytical results

obtained for hydrogen free DLC fabricated using the ICP system and highlights the

main findings for the RIBSD technique for DLC fabrication.

7.1 Main findings of the work on hydrogenated DLC films

The results for the ICP fabricated a-C:H films were presented and discussed in

Section 5.7, therefore in this section we provide a short summation and emphasise

the contributions of the research the field.

A range of a-C:H films were fabricated using the ICP system where a mixture

of Ar/CH4 gas was used as a constant intensity plasma source. Applied DC negative

bias voltage in the range of - 250 to - 400 V with 50 V increments was used to

produce films of different microstructure and properties. Nanoindentation

measurements, UV-Vis Raman spectroscopy, IR and XPS C1s core-level x-ray

photoelectron spectroscopy analysis were carried out on a-C:H film deposited onto Si

substrates. The films synthesised at higher negative substrate bias were found to

display superior mechanical properties (hardness and Young’s modulus) than those

at low bias. The ratio of sp3/sp2 hybridised fractions in all films was estimated by

deconvolution of XPS C1s core-level spectra and was found that the ratio is

essentially the same for all films fabricated and amounts to 28 ± 0.5 % (see Section

5.2)

.

112

Multi-wavelength Raman spectroscopy was used to examine sp2 and sp3 phases

of a-C:H samples. Raman scattering with visible light excitation, specifically 633

nm, 532 nm, enables examination of the sp2 component, whereas Raman in the UV

range (325 nm and 244 nm) can probe both the sp2 and sp3 phases. The Raman

results revealed (see Section 5.3) that the sp2 constituent in films fabricated at higher

bias displayed higher degree of bonding disorder in both aromatic rings and olefinic

chains; the rings were of smaller physical size and olefinic chains were significantly

shorter. It was also found that the amount of the sp3 phase was identical in all

fabricated films as indicated by the characteristic features (position, intensity,

FWHM) of the T peak of the UV Raman spectra. This is in agreement with the XPS

analysis (see Section 5.2). Inclusions of polyacetylene, the simplest of π- conjugated

polymers, were also found in all a-C:H films as evidenced by the peak positioned at

approximately 1200 cm-1 that appeared in Raman spectra through all wavelengths

studied. Preliminary measurements using Rayleigh light scattering (see Section

5.3.1) to monitor the relative density of the fabricated film samples was conducted

and it was found that the magnitude of the Rayleigh light was higher for films

displaying superior mechanical properties, that is films synthesised at higher bias.

The IR spectroscopy analysis performed showed that the contributions from the

primary sp3 group –CH3 decreased significantly in samples fabricated at higher bias,

whereas contributions from secondary sp3 =C-H and =CH2 olefinic groups

increased. There was also some increase in contributions from sp2 groups like from

C=C-H aromatic and from =C-H olefinic at higher bias. We consider these

transformations in the a-C:H film's morphology are caused by the changes in

bombardment energy of incoming C and H ions due to variation of bias voltage. The

ions energy is higher at higher bias and the changes in a-C:H morphology are related

to a de-hydrogenation (hydrogen abstraction) effect and an increase of local densities

in sp2 and sp3 phases (see Section 5.4).

The Tauc band gap and surface conduction band gap were determined for all

fabricated a-C:H samples (see Section 5.5). The Tauc gap is derived from

fundamental absorption spectra in the N-IR range and, the surface gap is obtained

from the SPS measurements. Due to the different nature of probing techniques, that

is N-IR is probing the material bulk while the STS measurements are surface

113

sensitive, it was found that the Tauc gap is inversely related to the surface conduction

gap over the range of fabrication conditions of a-C:H films in this work. This relation

was attributed to perturbed π - π* conduction gap of the sp2 fraction.

Research on this project contributed to the existing knowledge by providing the

evidence that, mechanical properties of hydrogenated DLC (a-C:H) films with minor

variations in hardness and Young’s modulus (in our case the variation in hardness

was ~3 GPa and for elastic modulus ~30 GPa), are not solely controlled by the

amount of the sp3 constituent but are determined by the structural arrangement of the

sp2 constituent (see Section 1.2.1 and 1.2.1.1). The formation mechanism of

hydrogenated DLC films and the role of C-H species during the film growth were

also examined (see Section 1.2.1.1 and 3.1) providing a better understanding for

DLC formation out of a hydrocarbon plasma medium. For an open plasma deposition

(ICP) method an increase in the negative bias applied to the substrate resulted in

formation of a film with higher fraction of olefinic groups both the sp3 and sp2 like.

Notably, the increase of bias over the relatively narrow range used did not contribute

to the increase in the sp3 fraction in the deposited films as observed in similar

deposition systems, but contributed to an increase in the structural disorder of the sp3

phase. This was indicated by reduced size and increased asymmetry of olefinic

domains. The sp2 fraction in all films become altered to a greater extent with

increasing negative bias, such as a decrease in the amount of sp2 aromatic sites and,

consequently an increase in the number of sp2 olefinic sites. The variation of bias

also affects the physical size and ordering of sp2 sites, they become more asymmetric

and smaller as the energy of C-H species increases (see Section 5.3 and 5.4).

It was shown that a common analytical technique, such as the XPS C1s core-

level analysis (see Section 4.3 and 5.2), when used in application to DLC

microstructure, is able to deliver consistent results determining the sp3/sp2 ratio,

when known fitting constraints (FWHM, peak separation, peak positions) are

employed throughout the spectral deconvolution process. It was demonstrated that

Raman spectroscopy (see Section 4.4 and 5.3), especially in the UV range, provides

comprehensive information about the microstructure of DLC films, including the

character of sp2 and sp3 hybridised bonding and their respective proportions. We

have also shown using provisional experiments that the magnitude of Rayleigh

114

scattered light is related to mechanical properties of DLC films such as the film

density (see Section 4.4.2).

The use of Pearson VII line function for spectral fitting of, for example, Raman

spectra was also examined, and we have presented theoretical background which

supports the suitability of this lineshape as compared to the currently used Breit–

Wigner–Fano line function (see Section 4.4.1).

Contributions from trans-polyacetylene (see Section 5.3) were also identified in

hydrogenated DLC as evidenced by the D* band indicated the complexity of DLC

microstructure.

Valuable information about the approach that allows unambiguous

determination of key mechanical parameters, such as hardness and Young’s modulus

from common nanoindentation data was also presented (see Section 5.1). The

measurements performed clearly indicate the deformation responses to an applied

load and, in addition, the use of the relative force propagation diagram was proposed

to supplement the extraction of information from the load propagation curves. The

relative ratio of hardness to Young's modulus was proposed to acquire additional

information about stress concentration in examined DLC samples. The

measurements performed showed that comprehensive information about micro-

mechanical properties of tested DLC films can be obtained from nanoindentation

measurements alone. The IR spectroscopy measurements (see Section 4.5 and 5.4)

provided valuable information about detailed morphological changes that take place

during DLC growth in the ICP reactor. We have monitored and analysed changes of

individual sp3 and sp2 bonding groups and co-related them with changes in C- and H-

bombardment energy during DLC growth process. The discrepancy between the

Tauc gap and the surface conduction gap was illustrated and explained (see Section

4.6 and 5.5). It is now known that the discrepancies are due to the differences of the

probing nature of IR and the SPS methods and are also related to the appearance of

the perturbed π - π* gap of sp2 phase.

115

7.2 Main findings of the work on development and

investigation of a new deposition technique (the RIBSD).

Hydrogen free DLC films were fabricated using the RIBSD deposition method

(see Section 1.2.1.2 and 3.2). The RIBSD system employed a single ion beam source

for fabrication of DLC films and both a sputter target (HOPG) and a substrate during

the sputtering arrangement were positioned in a close proximity to each other and

were tilted relative to each other. In this arrangement the incoming ions were

sputtering the target and were bombarding the growing film. The Ar and Xe ions

were used in experiments and low energy Kauffman type ion gun was operated

within 0.2 – 1.2 keV range.

It was found that the impinging bombardment was not favourable for growth of

DLC films as the bombardment caused secondary re-sputtering process that inhibited

sp3 formation on the surface of the substrate (see Section 3.2 and 6.3). That is when

the angle of the substrate, αs was 10° (impinging bombardment is present) only sp2

rich amorphous carbon films were fabricated. In such films the amount of the sp3

fraction was found to be very low, of less than 8 % as determined by the XPS C1s

measurements (see Section 4.3 and 6.2).

In the absence of impinging bombardment, that is when αs was 0° (substrate is

positioned parallel to the ion beam axis), DLC of low to medium sp3 content were

produced (see Section 3.2 and 6.3). The sp3 content was found to vary with the angle

of the target, αt to the ion beam axis. The values of αt studied were 15, 30 and 45°

(see Section 3.2). The sp3 content was found to be at maximum for all bombardment

energies studied for the αt of 30° and amounted to 43 % for Ar and 45 % for Xe

deposited films at maximum bombardment energy of 1.2 keV. The sp3 fraction was

principally determined by XPS C1s measurements (see Section 6.2). When no

impinging bombardment was present the increase of ion energy from 0.2 to 1.2 keV

contributed to significant structural changes in fabricated DLC films. The changes

were from predominantly sp2 graphitic-like bonding to tetrahedral sp3 bonding

arrangement.

116

The project contributed to the existing body of knowledge by developing an

original deposition technique (see Section 1.2.1.2 and 3.2) for fabrication of quality

DLC coatings that can be seamlessly integrated into nearly any vacuum facility and,

the RIBSD can be set up and operated in a thin film optical deposition chamber as

considered by the industrial partner of this investigation Laserdyne Pty Ltd. In fact,

any type of thin solid films can be fabricated using the RIBSD system provided that

selected target material can be sputtered using energetic ions within the operational

range of 0.2 to 1.2 keV.

We elucidated the effects of varying bombardment ion energy, ion types and

target and substrate geometry relative to the incoming ion beam onto the formation

of sp3 phase in DLC films (see Section 3.2 and 6.3).

We also found that the XPS C1s analysis (see Section 6.2) in application to

DLC provides reliable results determining the sp2/sp3 ratio as confirmed by the UV

Raman (325 nm) measurements (see Section 6.1 and 6.3). This is one of the most

important finding since there are very few reports that address the UV Raman (T

peak parameters) vs. XPS C1s compatibility.

7.3 Future outlook

The investigation of the formation mechanism and properties of hydrogenated

DLC films in the ICP reactor yielded a conclusion that mechanical properties in DLC

films are controlled by ordering of the sp2 phase in films where the sp3 content is

equal (see Section 5.7). In order to further prove this conclusion, additional analytical

measurements (see Section 4.1) will to be necessary, such as the methods that allow

a direct measurement of the amount of sp3 phase and the hydrogen content. Methods

like NEXAS, EELS, X-ray reflectivity or NMR measurements should suffice.

117

In order to further investigate capabilities of the RIBSD method for fabrication

of DLC films it will be useful to explore target bombardment with ions of much

higher energy than used in this project. The energies suggested for ion bombardment

will be in the range of 2.5 to 5 keV producing ejected C ions with energies close to

100 eV, however at this range of bombardment energies there will be progressive

stress accumulation in fabricated DLC films and measures should be taken to

suppress it. The increase of ion density would also lead to enhanced DLC film

formation. To produce hydrogenated DLC films, the RIBSD experiments can be

repeated with hydrocarbon gas as the ion source, or the experiments can be

conducted in atmosphere filled with a hydrocarbon gas and noble gas ions are used

for sputtering.

The authors performed the RIBSD experiments using N ions at the energies of

0.2 – 1.2 keV at the same experimental setting as used for Ar and Xe sputtering. The

results of these experiments are being analysed and will be reported elsewhere.

118

119

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