Maksym Rybachuk Thesis
Transcript of Maksym Rybachuk Thesis
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
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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
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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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