Chemical Vapor Deposition of Low-Dielectric Constant Organosilicon-Based Thin Films

A Thesis Presented

by

Nariné Razmik Malkhasyan

to

The Department of Chemical Engineering

In partial fulfillment of the requirements For the degree of

Master of Science

In

Chemical Engineering

Northeastern University Boston, Massachusetts

April 14, 2009

ii

ACKNOWLEDGEMENTS

I would like to express my gratitude to my advisor Prof. Albert Sacco, Jr. for his

continuous support and encouragement, as well as all his contribution to my scientific

understanding. I would like to acknowledge my advisor Prof. Daniel Burkey for his

academic as well as moral support, his valuable insight, and all the helpful advice. I want

to thank Prof. Katherine Ziemer for her guidance, enthusiasm and eagerness to help.

I am grateful to my lab mates at the Thin Film and CVD Laboratory: Courtney

Pfluger, Brian McMahon, Stephanie Fernandez, Anna Poehler, and Kyle Stephens, for

their help in the laboratory and the much needed occasional distraction from research.

I would like to thank Mike Dunlevy, who was my support system throughout my

years here at Northeastern. Thank you for helping me see the lighter side of things. All

life’s problems are easier with you by my side.

Lastly, I must thank my mother, Minna Gurgenyan, for her endless belief in my

abilities, her tireless encouragement, and for being the most important influence in my

life. I hope that some day I can be half the woman that you are.

iii

ABSTRACT

In contemporary semiconductor devices speed and functionality depends on

signal propagation and, subsequently, on the geometrical design of the device and its

material properties. Power dissipation, current leakage, and cross-talk noise are the

issues plaguing signal propagation. These issues can be crucial at the dimensions of

ultra-large scale integrated devices, as they increase exponentially with decreasing

dielectric layer thickness. Thus, having a dielectric with excellent insulating properties is

critical, since it can abate these issues. Organosilicon thin films are a natural choice for

dielectric layers, since they are similar to the previously used silicon dioxide, can be

deposited using the same technology, but have significantly lower dielectric constants.

Organosilicon films, in addition to being good electrical insulators, have been

shown to be biocompatible, which makes them promising candidates for passivating

coatings in bioimplantable devices, such as neuroprosthetics. The solvent-free chemical

vapor deposition technique insures that the biocompatibility of the passivation coating

will not be compromised by solvent leaching. Additionally, it allows the deposition of

the passivation coating in conjunction with a reactive polymer coating, which can then be

functionalized with biomolecules.

The goal of this investigation was to obtain an organosilicon thin film with lowest

possible dielectric constant via CVD: co-polymerizing a cyclic precursor with intrinsic

porosity and a linear precursor acting as a spacer for additional porosity.

Trivinyltrimethylcyclotrisiloxane (V3D3) was the cyclic precursor used for the

depositions. Vinylmethyldieothoxysilane (VDEMS), divinyldimethylsilane (S1),

iv

divinyltetramethyldisiloxane (S3), and divinylhexamethyltrisiloxane (S5) were the linear

spacer precursors used. This study also investigated depositions of V3D3 film with a top

layer of poly(acrylic acid) (pAA) that could be functionalized with biomolecules.

Plasma enhanced CVD and initiated CVD techniques were used for deposition of

V3D3 and V3D3-spacer films. PECVD was shown to be too aggressive as a deposition

technique, and resulted in precursor structure loss. The lowest as-deposit refractive index

obtained for a V3D3 film deposited at a low substrate temperature (25 ºC) and low power

(10 W) was 1.482 ± 0.001. By contrast the mildest iCVD conditions yielded a V3D3 film

with significant retention of cyclic structure, as evidenced by FTIR spectra and peak

deconvolution, and with a refractive index of 1.454 ± 0.002. This showed that iCVD was

a more suitable milder alternative for depositing films from precursors with delicate

functionality.

Both techniques yielded V3D3-spacer films with refractive indices either higher

or statistically the same than those of V3D3-only films deposited at the same conditions.

This suggested that the linear siloxane molecules were not acting as spacer molecules as

had been hypothesized. It was likely that the spacers polymerized through the vinyl

bonds only, similar to acrylates, resulting in densification rather than formation of

additional porosity. Additionally, in PECVD the abundance of energy could have

resulted in precursor fragmentation leading to film densification, as opposed to additional

pore formation.

The ratio of the monomer partial pressure to its saturation pressure (PM/Psat) was

proven to be a crucial variable in iCVD by Lau and Gleason [52, 53], particularly for co-

polymerizations, since it quantified the amount of monomer adsorbed onto the substrate,

v

which, in turn, controlled the deposition rate. It was shown that a V3D3-S5 film (film

that added divinylhexamethyltrisiloxane to V3D3) that was deposited with equal PM/Psat

ratios for both monomers had a refractive index of 1.453 ± 0.010, whereas the film

deposited with different PM/Psat values had a refractive index of 1.461 ± 0.002.

V3D3 films were investigated as base layers for biofunctionalized films.

Biofunctionalization would be achieved by protein attachment to carboxylic acid groups

in the top pAA layer, which had to be stable in aqueous environment. It was shown that

increasing plasma power and increasing initiator flowrate improved the film stability in

water, but resulted in the decrease of the carbonyl peak in their FTIR spectra, suggesting

that cross-linking took place at the expense of the carboxyl acid functionality. This was

hypothesized to be due to the fact that the abundance of tBPO radicals in the plasma not

only initialized the linear polymerization of AA through the vinyl chemistry but also

reacted with other available bonds in the growing pAA linear chains and cross-linked

these, thus, consuming carboxylic acid functional groups. Co-polymerization with

ethylene glycol diacrylate (EGDA) showed to have the desirable effect on film stability

without depleting the films of their important functionality. A V3D3/pAA-co-EGDA

film was shown to have good thickness retention, ~95 %, in DI water for an hour, due to

cross-linking of AA with EGDA and the likely grafting between the V3D3 film and the

AA/EGDA top layer. Protein tethering on the same combination V3D3/pAA-co-EGDA

film using fluorescein isothiocyanate anti-mouse immunoglobulin G was shown to be

successful by fluorescent microscopy proving that was is possible to combine electrically

insulating and biocompatible and biofunctional surfaces into one film using CVD

techniques.

vi

TABLE OF CONTENTS

LIST OF FIGURES viii

LIST OF TABLES xv

1. INTRODUCTION 1

1.1 Interconnect Materials 2

1.2 Biocompatible Insulating Materials 13

2. CRITICAL LITERATURE REVIEW 17

2.1 Spin-on Technique 18

2.2 Sacrificial Porogen Technique 23

2.3 Intrinsic Porosity Technique 31

2.4 Molecular Architecture Technique 39

2.5 Initiated Chemical Vapor Deposition 43

2.6 Organosilicon Films as Biocompatible Insulators 53

2.7 Characterization Techniques 57

2.7.1 Fourier Transform Infrared Spectroscopy 58

2.7.2 Spectroscopic Ellipsometry 65

3. EXPERIMENTAL APPROACH 68

4. RESULTS AND DISCUSSION 78

4.1 Plasma Enhanced CVD 78

4.1.1 Vinylmethyldiethoxysilane Depositions 79

4.1.2 Trivinyltrimethylcyclotrisiloxane Depositions 91

vii

4.1.3 Trivinyltrimethylcyclotrisiloxane-Vinylmethyldiethoxysilane

Depositions 104

4.1.4 Trivinyltrimethylcyclotrisiloxane-Spacer Depositions 109

4.2 Initiated CVD 117

4.3 Deposition of Biocompatible Insulating Films 136

5. CONCLUSIONS 150

6. FUTURE WORK 160

7. NOMENCLATURE 161

8. ABBREVIATIONS 163

9. REFERENCES 165

viii

LIST OF FIGURES

Figure 1: A schematic of (a) tetrahedral silica unit, (b) an SSQ cage with eight

Si atoms (T8), and (c) an SSQ ladder structure given by Baklanov and

Maex [1] 5

Figure 2: Correlation between dielectric constants and porosities of the films

prepared by Ting et al [39] using different templates 20

Figure 3: Correlation between the mechanical properties: (a) Young’s modulus (E)

and (b) hardness (H) and dielectric constant of the low dielectric constant

films evaluated by Hijioka et al [36] using different templates 22

Figure 4: (a) Dielectric constant and (b) thickness change during annealing as a

function of the organic porogen fraction in the gas feed during deposition

of low dielectric constant films by Grill [4] 25

Figure 5: FTIR spectra of as-deposited organosiloxane films deposited at substrate

temperatures of 100, 200 and 300 ºC by Gates et al [42] 28

Figure 6: Extent of porogen decomposition as a function of annealing temperature

tracked by (a) comparison of C=O peak in the FTIR spectra and (b)

comparison of the normalized area under the C=O peak in films deposited

by Burkey and Gleason [41] 30

Figure 7: (a) Deposition rate and (b) dielectric constant of films deposited by Tada

et al [16] as a function of the partial pressure of cyclic siloxane monomer 33

Figure 8: Dielectric constant of 3V3RL films deposited by Tada et al [16] as a

function of (a) plasma power and (b) monomer partial pressure 34

ix

Figure 9: The Si(CH3)x stretching region of the FTIR spectrum for (a) TVMS, (b)

DVDMS, (c) V3D3, and (d) H4D4 films deposited by Burkey and

Gleason [14] 35

Figure 10: FTIR spectra of films deposited from V3D3/H2O at three different duty

cycles by Burkey and Gleason [15] 37

Figure 11: FTIR spectra of (a) D3 PECVD film deposited under continuous-wave

excitation and (b) D3 HFCVD film deposited at 1000 ºC by Pryce Lewis

et al [50] 45

Figure 12: FTIR spectra of (L) D3 films deposited at (a) 860, (b) 1000, and (c)

1100 ºC and of (R) D4 films deposited at (a) 800, (b) 900, and (c) 1000

ºC by Pryce Lewis et al [50] 46

Figure 13: FTIR spectra of the V3D3 monomer and the resulting iCVD polymer

deposited by O’Shaughnessy et al [26] 48

Figure 14: Deposition rate data for polymer growth as a function of filament

temperature for V3D3 films deposited by O’Shaughnessy et al [26] 49

Figure 15: Reaction mechanism proposed for iCVD polymerization by Lau and

Gleason [52] 50

Figure 16: Effect of monomer-saturated vapor pressure on (a) polymer deposition

rate and (b) number-average molecular weight for iCVD poly(alkyl

acrylate) films deposited by Lau and Gleason [52] 51

Figure 17: Effect of substrate temperature on polymer deposition rate of iCVD

ethyl acrylate films deposited by Lau and Gleason [52] at filament

temperature of () 285, () 310, and () 360 ºC 52

x

Figure 18: PC12 neuron growth in the presence of glass substrates both uncoated

and coated with iCVD V3D3 by O’Shaughnessy [8] 55

Figure 19: An FTIR spectrum typical of organosiloxane films 61

Figure 20: FTIR absorbance peak indicative of the degree of oxidation of silicon

atoms 63

Figure 21: Incorporation of thermally labile porogen in a bulk matrix 69

Figure 22: Cyclosiloxane precursors with vinyl and methyl side chain chemistry 70

Figure 23: Linear spacer molecules with different functional groups 71

Figure 24: Schematic representation of organosilicon films with molecular (due

to V3D3) and tunable (due to spacer molecules) porosity 72

Figure 25: Molecules used in deposition of biofunctionalizable coating 73

Figure 26: Photograph of the filament holder used in iCVD depositions 75

Figure 27: Overall FTIR spectra of the VDEMS films deposited at the conditions

specified by the design of the experiment matrix 81

Figure 28: Segment of FTIR spectra of VDEMS films deposited at the conditions

specified by the design of the experiment matrix 83

Figure 29: Segment of FTIR spectra of as-deposit VDEMS films deposited at low

substrate temperature 84

Figure 30: Segment of FTIR spectra of post-anneal VDEMS films deposited at low

substrate temperature 85

Figure 31: Segment of FTIR spectra of as-deposit VDEMS films deposited at

high substrate temperature 86

xi

Figure 32: Segment of FTIR spectra of post-anneal VDEMS films deposited at

high substrate temperature 87

Figure 33: (a) Siloxane and (b) organic ranges of FTIR spectra of as-deposit and

post-anneal VDEMS film deposited using 1 sccm VDEMS flowrate,

100 W plasma power, and 25 °C substrate temperature 88

Figure 34: (a) Siloxane and (b) organic ranges of FTIR spectra of as-deposit and

post-anneal VDEMS film deposited using 5 sccm VDEMS flowrate,

20 W plasma power, and 25 °C substrate temperature 90

Figure 35: Segment of FTIR spectrum of PECVD V3D3 film deposited at mild

plasma conditions 93

Figure 36: Deconvolution of the siloxane peak of the PECVD V3D3 film spectrum 94

Figure 37: FTIR spectra of the PECVD V3D3 film, deposited at mild plasma

conditions, before and after annealing. 95

Figure 38: Segment of FTIR spectra of the PECVD V3D3 film, deposited at mild

plasma conditions, before and after annealing 96

Figure 39: Deconvolution of the siloxane peak of the post-anneal PECVD V3D3

film spectrum 98

Figure 40: Segment of FTIR spectra of the PECVD V3D3 film, deposited at high

substrate temperature, before and after annealing 99

Figure 41: Deconvolution of the siloxane peak of the as-deposit PECVD V3D3 film

deposited at high substrate temperature 101

Figure 42: Deconvolution of the siloxane peak of the post-anneal PECVD V3D3

film deposited at high substrate temperature 102

xii

Figure 43: Comparison of FTIR spectra of V3D3-only films deposited at low

(25 °C) and high (200 °C) substrate temperatures 103

Figure 44: Comparison of PECVD V3D3-only and V3D3-VDEMS film spectra,

deposited at the same conditions 106

Figure 45: Deconvolution of the siloxane peak of the as-deposit PECVD V3D3–

VDEMS film deposited at high substrate temperature 108

Figure 46: Deconvolution of the siloxane peak of the post-anneal PECVD V3D3-

VDEMS film deposited at high substrate temperature 109

Figure 47: Segment of FTIR spectra of the PECVD V3D3-S3 films, deposited at

25 °C, 10 W and 200°C, 10 W 112

Figure 48: Comparison of V3D3-only and V3D3-spacer film spectra deposited at

200 °C and 10 W 113

Figure 49: Deconvolution of the siloxane peak of the as-deposit PECVD V3D3–S3

film deposited at 200 °C and 10 W 115

Figure 50: Deconvolution of the siloxane peak of the post-anneal PECVD V3D3-

S3 film deposited at 200 °C and 10 W 116

Figure 51: Segment of FTIR spectra of iCVD V3D3-only films deposited at

filament temperatures of 200 and 350 °C 119

Figure 52: Siloxane peak deconvolution of the as-deposit iCVD V3D3-only film

deposited at the filament temperature of 200 °C 120

Figure 53: Siloxane peak deconvolution of the as-deposit iCVD V3D3-only film

deposited at the filament temperature of 350 °C 121

xiii

Figure 54: Comparison of FTIR spectrum segments of PECVD and iCVD V3D3-

only films 123

Figure 55: Comparison of FTIR spectrum segments of PECVD and iCVD V3D3-

S3 films deposited with minimal energy input 126

Figure 56: Comparison of FTIR spectrum segments of PECVD and iCVD V3D3-

S5 films deposited with minimal energy input 127

Figure 57: iCVD V3D3-spacer films using S1, S3 and S5 deposited at the same

conditions 128

Figure 58: iCVD V3D3-only films using different Pm/Psat ratios controlled by

changing the substrate temperature 131

Figure 59: iCVD V3D3 and V3D3+S5 filpms deosited at the same conditions

with the same Pm/Psat ratios 133

Figure 60: Comparison of FTIR spectrum segments of PECVD (low substrate

temperature and low power) and iCVD (PM/Psat=0.42) V3D3-S5 films 135

Figure 61: Segment of FTIR spectra of pAA films deposited at varying plasma

powers 137

Figure 62: Increasing film thickness retention with increasing plasma powers

suggests that at higher powers precursor fractionation in plasma

yields more densely cross-linked AA films 138

Figure 63: Improvement in film thickness retention with the increase of tBPO

flowrate, which points to tBPO contributing to film cross-linking 139

Figure 64: Segment of FTIR spectra of pAA films deposited with varying

amounts if tBPO 140

xiv

Figure 65: Illustration of EGDA molecule used as a cross-linker between two

polymer chains [62] 141

Figure 66: Segment of FTIR spectra of pAA-co-EGDA films deposited with

varying EGDA temperatures 142

Figure 67: Increasing film thickness retention with increasing EGDA precursor

temperature points to EGDA acting as a cross-linker in the films 143

Figure 68: FTIR spectrum of V3D3/pAA-co-EGDA film 144

Figure 69: Siloxane range in the FTIR spectrum of V3D3/pAA-co-EGDA film 146

Figure 70: Carbonyl peak in the FTIR spectrum of V3D3/pAA-co-EGDA film 147

Figure 71: Fluorescent micrograph of FITC anti-mouse IgG protein tethered to

V3D3/AA film surface 149

xv

LIST OF TABLES

Table 1: Hardness (H) and Young’s modulus (E) of each porous film prepared by

Ting et al. [39] using different templates 20

Table 2: Comparison between electrical and thermal properties of films deposited

from DVTMDSO and HMDSO by Milella et al. [47] 40

Table 3: FTIR peak assignments for organosiloxane films 64

Table 4: List of chemicals used 74

Table 5: Orthogonal design of experiment matrix 80

Table 6: Optical properties and thermal stability data for the VDEMS films

deposited at the conditions specified by the design of the experiment

matrix 87

Table 7: Optical properties and thermal stability data for the V3D3-spacer films

deposited at varying substrate temperatures and deposition powers 110

Table 8: Optical properties for the V3D3-spacer films deposited using the iCVD

technique 124

Table 9: Comparison of optical properties of the V3D3-spacer films deposited

using the iCVD and PECVD techniques 124

Table 10: Comparison of deposition rates and optical properties of iCVD V3D3-

only films deposited using different PM/Psat ratios 130

1

1. Introduction

Organosilicon films are amorphous carbon-doped glass materials that are known

for their excellent properties as electrical insulators [1-4] and biocompatible coatings [5-

8]. These properties stem from the molecular structure of these films: the incorporated

organic groups reduce the polarizability of the films due to their non-ionic nature and are

often present in the form of end-capping methyl groups, which create void spaces around

themselves and thus decrease film density, both of which result in a decrease of the

dielectric constant of film.

The dielectric constant of a material, k, is described by the Clausius-Mossotti

equation [1]:

( )( ) απ Nkk

34

21

=+− (1)

0εε=k (2)

Where ε is the relative permittivity of the material, ε0 is the permittivity of

vacuum, N is the number of molecules per unit volume (density), and α is the total

polarizability. Hence the dielectric constant of a material can be controlled by controlling

its density and polarizability. Polarizability is a measure of tendency of charge

distribution to be distorted from its original shape by an external electric field. Reducing

the number of ionic bonds in the material, reducing the electron density, i.e. introducing

smaller elements, as well as incorporating atoms with high electronegativity, which

would strongly attract electrons towards themselves, are all methods to reduce

polarizability.

2

In addition, reducing film density can yield dielectric constants significantly

lower than reducing film polarizability alone. An important way to reduce film density is

the introduction of pores. The dielectric constant of a porous film then depends on the

porosity and the dielectric constant of the film skeleton, as well as that of the material

inside the pores:

( )( ) ( ) ( )

( )21

121

21

1

1

+−

−++−

=+−

s

s

r

r

kk

PkkP

kk (3)

Where kr is the dielectric constant of the porous material, k1 is the dielectric

constant of the material inside the pores (unity for air), ks is the dielectric constant of the

film skeleton, and P is the porosity of the material [1]. Thus, a film with a low dielectric

constant can be obtained by using skeleton materials with low dielectric constants and

then by introducing pores into the film structure, to reduce the dielectric constant even

further.

Organosilicon thin films have found use as dielectric interconnect materials in

semiconductor devices due to the ease of their incorporation into existing processes, since

they are similar to silicon dioxide: the interconnect material used previously. Along with

being good electrical insulators, these thin films have also been shown to be

biocompatible [5, 6, 54], which makes them a logical choice for coatings for electrical

devices that need to perform in biological environments such as neuroprosthetics.

1.1

The semiconductor industry was a 175 billion dollar industry in 2005 [9] and

continues to grow steadily. Due to the consumer demand for smaller, sleeker, more

portable, and more integrated electronics the functionality of a single chip has had to

Interconnect Materials

3

increase while the size of the chip grew smaller. According to Moore’s Law, set forth by

Intel executive Gordon Moore [9], the market demand for functionality per chip doubles

every 1.5 to 2 years. And as Moore’s Law has proven to be a successful macro-trend for

leading-edge semiconductor products for the past 30 years, the industry has strained to

keep up.

Modern ultra-large scale integration (ULSI) devices contain 108-109 transistors in

an area smaller than 1 cm2 [9]. The basic elements in the device need to be

interconnected for the device to be functional. As the device size shrinks the level of

interconnection and the length of the metal wiring increase. Since the functionality and

speed of the device depends on signal propagation between the elements, which, in turn,

is dependent on the properties of the metal wiring as well as on those of the inter-level

dielectric material. Resistive-capacitive (RC) delay, which hinders the signal speed in the

device and is a function of the resistance of the metal and the capacitance of the dielectric

material, is given by:

+= 2

2

2

2

042

TL

PLkRC ερ (4)

Where ρ is the metal resistivity, ε0 is vacuum permittivity, k is the relative

dielectric constant of the interlayer dielectric, P is the pitch, the separation between the

metal lines, T is the metal thickness and L is the metal line length [1]. Since electrical

resistance is the reciprocal of conductance, metal interconnects with higher resistivity can

be expected to be a poor choice for circuit manufacturing. Also, the higher the

capacitance of the inter-layer dielectric material, the better its ability to store energy from

the electric field applied by the metal interconnects, thus diminishing the electrical charge

passing through the circuit wires. Based on Equation 4, if the geometry of a device is the

4

same, the signal delay in the device can be reduced by lowering either the resistivity of

the metal or the dielectric constant of the interconnect material. Since copper is now the

most common metal used in microelectronics, the interconnect material is the

determining factor for signal delay [2]. Power consumption is another major concern

since increasing frequencies and higher densities result in an increase in power

consumption. Leakage current, mobile charge carriers tunneling through the insulating

layer, contributes to an increase in power consumption, and can be influenced

significantly by the dielectric constant of the interconnect material [2]. The lower the

dielectric constant of the insulating layer, the more difficult it is for an electron to pass

through it. This is crucial at the dimensions of ULSI devices, as leakage current

increases exponentially with decreasing thickness of the dielectric layer.

Historically, silicon dioxide, which has a dielectric constant, k, of about 4.1, has

been used as interconnect dielectric material. Using a material with a lower dielectric

constant would aid in reduction of signal propagation delay, cross-talk noise, etc.

Presently, several different types of materials are used for preparing low dielectric

constant films: silica-based materials, silsesquioxanes (SSQ), and organic polymers [1,

2]. Silica-based and SSQ materials have properties similar to those of silicon oxide,

making their integration into existing technology easier. Silica-based materials have a

typical SiO2 tetrahedral structure, but some of the oxygen molecules are substituted with

a fluorine or a –CH3 group to lower film polarizability and density, since, for example, an

end-capping methyl group creates void space around itself. These types of films are

mainly prepared via the chemical vapor deposition (CVD) method. SSQ materials are

organic polymers with the empirical formula ( )n

SiOR 23− , where R can be a hydrogen

5

atom, or an alkyl, alkenyl, alcoxy, or an aryl group [2]. In SSQ materials Si and O atoms

are arranged in a cage or a ladder structure, shown in Figure 1, which creates free volume

decreasing film density. This low density and the lower polarizability of Si-R bonds both

contribute to the low dielectric constant of SSQ materials. These types of films are

generally prepared by the spin-on polymerization technology. Organic polymer films are

typically non-polar in order to keep the dielectric constant low. Impurities in the film or

unsaturated terminal groups contribute to the polarizability of such films. Organic

polymer materials are more temperature sensitive than silica-based or SSQ films making

their integration more difficult.

Figure 1: A schematic of (a) tetrahedral silica unit, (b) an SSQ cage with eight Si atoms (T8), and (c) an SSQ ladder structure given by Baklanov and Maex [1]. Si atoms are shown in black, O atoms in gray.

As previously mentioned, SSQ films are deposited by spin-on polymerization, for

which the dielectric precursor should be available in “sol” form: primary particles

6

dispersed in a solvent. This liquid sol is dispensed onto the substrate, which is then

placed on a spinner. The thickness of the film is the result of a balance between

centrifugal forces and viscous forces. The spinning step is followed by a bake step at

about 250 ºC to remove the solvent and then by a curing step at 350-600 ºC to induce

final cross-linking of the polymer chains. As with the other types of dielectric films,

porosity is one of the main ways to reduce the dielectric constant of spin-on films.

Porosity can be introduced into spin-on films through sol-gel processes [35, 39, 40] or via

processes that utilize sacrificial porogens [11, 38]. There are two sol-gel methods that

can be used: one based on aging processes and another based on self-assembly of the

primary particles in the sol [35, 39]. In order to prevent the pores from collapsing during

the baking step, it is important to ensure that a rigid skeleton structure is in place before

extraction of the solvent. The aging step, which is done by changing the pH of the

solvent, basically accelerates the sol-gel reactions prior to the baking step. But the

technique based on aging processes does not allow for independent control of pore size

and total porosity, as the two are related in this technique: greater porosity means larger

pore size. However, the level of porosity can be tuned by changing the solvent to solid

ratio [1, 2]. The other sol-gel technique for obtaining subtractive porosity in spin-on

films is based on the organized aggregation of the primary particles. In this case the pore

size and total porosity of the film are related to the primary way the particles are ordered

[35, 39]. The technique of obtaining subtractive porosity using sacrificial porogen is

based on the addition of porogen particles into the sol, which would withstand the bake

step but would volatilize and be removed during the curing step. Ideally the porosity of

the final film would be related to the amount of porogen in the sol and the pore size

7

would be related to the size of the sacrificial porogen. For this ideal scenario the porogen

has to be evenly dispersed in the sol and has to be compatible with the matrix material to

avoid phase separation. The porogen method allows for better cross-linking of the matrix

at the time when pores are created compared with the sol-gel methods.

Silica-based dielectric films are mainly obtained via CVD, during which

precursor species are created from volatile source compounds. In case of plasma

enhanced CVD (PECVD) electron-impact dissociation and ionization are the steps

leading to film deposition [10]. The mechanism of plasma polymerization is very

different from that of conventional polymerization: monofunctional compounds can be

polymerized with the aid of plasma, and also the high rate of initiation leads to a high

concentration of reactive species, and chain-termination reactions dominate over

propagation reactions. The products formed in the termination step are reinitiated and

continue to polymerize, yielding a highly cross-linked and irregular network. Most low

dielectric films obtained via CVD are doped versions of silicon oxide, with main dopants

being fluorine atoms or alkyl groups. The addition of fluorine lowers the polarizability of

the films, whereas the addition of alkyl groups also lowers the polarizability of the film as

well as decreases its density due to their large molecular volume. Additional porosity can

be obtained when thermally unstable Si-F or Si-alkyl groups are removed during post-

deposit high temperature steps, although thermal annealing of F-doped films results in

their densification [2].

Currently CVD is the preferable method for deposition of low dielectric constant

interconnect films in industry due to productivity and cost factors. CVD provides an

uninterrupted process sequence, since there are many other chip processing steps that also

8

require CVD-based steps, such as deposition of a barrier layer, etc. Spin-on technique

disrupts this sequence and would require spin-on coaters and furnaces, meaning a greater

capital investment [3]. Another factor that makes CVD silica-based, also known as

organosilicon, films more preferable is their mechanical properties: spin-on dielectric

films develop cracks when they reach a certain thickness [11]. Due to their more random

3-dimensional cross-linked structure CVD organosilicon films are not as prone to stress

cracking as are the spin-on polymeric films and can be used as much thicker films in

interconnect structures [4, 11].

There are numerous parameters in PECVD that can be controlled in order to tune

the final qualities of the deposited film, such as the power and frequency of the electrical

source, total and partial pressures in the reactor, gas flowrates into the reactor, pumping

speed, substrate temperature, and reactor design factors like overall geometry [10]. The

increase of power results in increased density of energetic electrons in the reactor, the

abundance of which, in turn, would lead to even greater monomer initiation and

fragmentation [10, 12]. The increase of substrate temperature has an effect partly similar

to that of increased power, though a different mechanism is involved: the polymer

structure is affected by continual pyrolysis of the film being deposited as well as by the

surface mobility of the precursor species on the substrate. An increase in either power or

substrate temperature has been observed to cause the elimination of carbon and hydrogen

containing groups and to yield a highly cross-linked mainly inorganic film [10]. On the

other hand, the deposition rate of the film increases with an increase in power due to

faster rate of monomer initiation, but it decreases with an increase in substrate

temperature, since it causes the sticking coefficient of the impinging precursor species to

9

drop. Deposition rate can typically be expected to increase with increasing flowrate in

the absence of mass transfer limitations: increasing the flowrate means that more

monomer species are made available for initiation. But an increase in flowrate leads to a

decrease in residence time of precursors in the reactor, thus decreasing the degree of

interaction between the species. This may be desirable if the goal is to keep the monomer

structure intact as the monomer gets more fractionated the longer it is present in the

plasma. On the other hand, extremely short residence times result in little or no film

being deposited since the precursors do not have time to adhere to the substrate. This can

be counteracted by increasing the deposition power. The pressure inside the reactor has

several effects, one of which is the effect on residence time, which is directly

proportional to the total pressure. Pressure also affects the mean free path of the

precursor species inside the reactor, since higher pressure implies higher precursor

density. Thus polymerization occurs more readily at higher pressures as there are more

reactive species present. Decreasing the partial pressure of the monomer in the reactor at

high flowrates allows to preserve the molecular structure of the monomer [12].

The most important properties of low dielectric constant organosilicon films are

related to their pore structure and porosity. The maximum pore size must be sufficiently

smaller than the minimum feature size of the device to provide structural stability for the

subsequent metallization layer of the device [13]. There are two categories of film

porosity: intrinsic and induced. Intrinsic porosity refers to molecular-size porosity

already built into the film structure, i.e. preserved ring structure of a cyclic monomer.

Induced porosity can be achieved via copolymerization of the film matrix with thermally

sensitive sacrificial porogen molecules that are further removed post-deposition through a

10

curing step or via enhancement of the molecular architecture through the addition of

spacer molecules. In most spin-on films the porogen species maintain their original

structure, whereas during PECVD both the matrix and porogen precursors are fragmented

into reactive species. Thus, for spin-on films the choice of chemistry and curing process

are crucial, as opposed to process parameters that control precursor fragmentation and

curing conditions that are important for obtaining this type of induced porosity in

PECVD films. Obtaining induced porosity utilizing sacrificial porogen molecules

increases the complexity of film deposition because of interplay between the main matrix

precursor and additional chemistry used as porogen. Complexity of deposition causes

difficulty controlling final films properties such as porosity, hardness, etc. Hardness and

elastic modulus of a film are important as indicators whether the film will survive the

integration process into a device: integration steps such as chemical-mechanical polishing

of damascene copper interconnects put significant amounts of mechanical stress onto the

structures, and poor mechanical properties of the dielectric layers can compromise the

reliability of the device. It has been reported that in order to withstand the mechanical

stresses associated with industrial processing steps dielectric films need a hardness of

more than 0.5 GPa [14]. To create a rigid matrix high power oxidant rich plasma is often

needed [14-17]. On the other hand, it is important to preserve the chemical structure of

the sacrificial porogen moieties so as to retain their volatile qualities, and high plasma

power and oxygen content can result in alteration of their chemical structure [15]. Curing

the film to volatilize the porogen species and generate pores in the film can be done by

providing thermal energy or photon energy (UV cure). Curing temperature is typically

limited to ≤400 ºC, which sometimes can be insufficient for removin g all of porogen

11

species from the film. One way to supply additional energy is to use a UV light source in

addition to the thermal cure, which will help to enhance the kinetics of the removal of the

porogen species without compromising the mechanical structure of the film by applying

excessive heat. Cure pressure can also play a significant part, as it has been shown that

lower pressure accelerates the removal of volatile species [18].

Porous films are classified based on their pore size: mesoporous (pore diameter >

2 nm), microporous (pore diameter < 2 nm), films with bi-modal porosity (containing

both micro- and meso- pores), films with ordered periodical pores, and films with

embedded voids interconnected by micropores. Pore connectivity is an important

characteristic since it can cause problems during the integration of the film into the

device: the material may present problems with diffusion of chemical substances used in

later processing steps, such as ones used during deposition of diffusion barrier or copper

interconnect layer, during chemical-mechanical polishing of copper, or water adsorption.

While the greater the degree of porosity in the film the lower its dielectric constant is, it

also means greater degree and depth of plasma damage as well as greater the penetration

of precursors during barrier film deposition. This can cause issues with film

hydrophobicity, which is crucial, as even small amounts of adsorbed water (dielectric

constant of 80), would significantly increase the dielectric constant value of the film.

Hydrophobicity of the film is typically achieved via introduction of H, CHx, and other

organic groups, but Si-H, Si-C, and Si-CH3 bonds can easily break when exposed to

plasmas, such as oxygen plasma used for etching or stripping, thus making the film less

hydrophobic. The depth of this plasma damage is directly related to film porosity and

pore interconnectivity [19]. Also a material that it too porous presents problems for the

12

adhesion of the barrier layer, which needs to be dense yet thin to prevent copper

penetration into the porous dielectric without making a significant impact on the

dielectric constant of the insulating layer between the copper wires. Porosity of the film

also has an effect on its thermal and mechanical properties. It is crucial for low dielectric

constant films to be able to withstand the high temperatures during further device

manufacturing steps that can reach upwards of 400 ºC. Thermal conductivity of porous

low dielectric constant films is yet another important criterion: with more features

concentrated on small devices, temperature-induced reliability issues can arise. The heat

conduction in disordered dielectrics may be considered a propagation of anharmonic

elastic waves through a continuum. This propagation occurs via interaction between the

quanta of thermal phonons, which at temperatures above 50 K can be considered to be a

diffusion process, and thus dependent on film porosity [13, 19]. The dielectric film also

needs to withstand the mechanical stress during the excess copper removal polish, after

the copper layer is deposited onto the dielectric layer. Mechanical properties quickly

deteriorate as film porosity increases: Young’s modulus, also known as elastic modulus,

which is a measure of stiffness of the material, decreases from 76 GPa for bulk SiO2 to

less than 10 GPa for SSQ films with 50% porosity [19]. Mechanical and thermal

properties of porous films appear to change similarly with change in porosity. This is yet

another aspect that makes CVD organosilicon films preferable to spin-on films: CVD

films have been reported to have pore sizes significantly smaller than those observed in

porous spin-on films of similar dielectric constants [10, 11].

13

1.2

Many industries have become dependent on advances in biotechnology in recent

years: from agriculture and food processing, to all aspects of medicine. The development

of cutting-edge biomedical applications, such as neuroprosthetic devices, has created a

need to interface sophisticated devices with biological environments [20]. Due to this

need a whole new line of research has arisen, one concerned with methodologies of

obtaining biocompatible coatings in order to functionalize biomaterial interfaces, which,

in turn, would have a great impact on the control of cell and tissue responses in vitro and

in vivo [20-22, 27]. Devices, such as implants, that come in close contact with biological

media need to be functionalized so as to minimize protein adsorption and cellular

interactions to avoid inflammation and fibrous encapsulation. On the other hand,

interactions between surfaces and biomolecules are crucial for biosensors and other

bioanalytical in vivo applications [23, 27]. One of the main problems with implant

therapy is chronic inflammation, which is the body’s reaction to a foreign object. This

inflammation can be caused by leaking of residual monomer or solvent from the polymer

implant into the tissue, or as a direct response to the nature of the implant surface [24-26].

A person can experience severe fevers and septic shock, which would make the removal

of the implant imperative. As a consequence, implants are built in their entirety, or, at the

very least, are coated with plastic materials that would minimize chronic inflammation.

Hence, in order to improve biocompatibility, biomaterials research has focused on

functionalization of material surfaces to control the biological response of the host.

Biocompatible Insulating Materials

Chemical vapor deposition is a preferred method for modifying surfaces for

biological applications, since the biocompatibility of materials obtained via the traditional

14

solution polymerization route can be compromised by leaching of either leftover solvent

or unreacted monomers [24-26]. CVD is an all-dry process, which eliminates the need

for solvents that could cause inflammation and tissue damage. All residual monomer is

also removed from the resulting material by vacuum degassing after deposition. CVD

provides the advantages of the ability to deposit different precursors together or

sequentially, the great variety of possible chemical structures obtained, as well as the

ability to coat two- and three-dimensional, as well as porous substrates [27].

Additionally, CVD allows exact control over film thickness, retention of monomer

functionality, modification of the very first layers of the surface while keeping the bulk

material properties [20, 22, 25]. Since CVD films are conformal and do not have wetting

issues associated with surface tension that can plague solution-phase polymerization, they

can be used to coat structures with complex topology and small dimensions, such as

neurological electrodes [26, 28]. The most widely-used thin films for biomedical devices

are probably the non-functionalized poly(p-xylylene) coatings, that have been

commercialized under the parylene brand [20]. Commonly, these are obtained through

CVD polymerization of [2.2]paracyclophanes (the so-called Gorham process) [29] and

are used in a wide range of applications, such as stents, cardiac pacemakers and

defibrillators, neuronal probes, orthodontic devices, catheters, electrosurgical tools, and

bioMEMS applications [30].

CVD thin films and organosiloxane films in particular are of significant interest in

the field of neuroprosthetics for the same reason as in the field of dielectrics: their

properties as an electrical insulator. Neuroprosthetic devices are interfaced with the

nervous system and act as a substitute for lost functions such as vision, hearing or

15

movement. For example, neuroprosthetic devices utilize the firing patterns of single

neurons to control electronic devices external to the subject. For optimal performance

neuroprosthetic devices have to be biocompatible and bioresistive. That is, they not only

have to be accepted by the biological system they are implanted into, but also have to be

chemically and physically stable in the biological environment [31]. Additionally, it is

important to minimize electrical noise in neuroprosthetic devices in order to ensure that

the signal received is from one neuron only, and to increase signal sensitivity. This can

be achieved by coating the device with a passivation layer with high electrical resistance.

Defined and stable surface properties along with the capability for biomolecule

attachment onto the surface are the most important features of biodevices [27, 32]. A

silicon dioxide layer can be used as an interlayer between a silicon-based device and the

functional coating. Reactive polymer coatings can improve the interfacial

biocompatibility of biomedical devices as they represent a designable interlayer.

Aldehyde, anhydride, or active ester groups allow for protein attachment; amino groups

can be used to control surface charges; while alkyl or fluoroalkyl groups provide

hydrophobic interfaces for electrochromatographic applications [20]. Carboxyl

functional groups, such as those in poly(acrylic acid) can serve as excellent anchor points

for immobilization of bioactive species and protein adsorption [23, 34]. Functional

coatings used to modify surfaces of biodevices include physisorbed and chemisorbed

materials as well as coating agents covalently bound to substrates. Adsorbed coatings are

less stable but easier to prepare, while chemical surface modification can be quite labor

intensive. For example, the hydrophobicity and chemical resistance of one of the most

common substrate materials in microfluidic devices, poly(dimethylsiloxane) (PDMS), are

16

often in need of modification to improve the biocompatibility of the device. The surface

modification methods include UV-induced polymer grafting, plasma polymerization, and

atom transfer radical polymerization [32], and a few of the techniques replace the methyl

groups of PDMS with hydroxyl groups to render the surface more hydrophilic. Most

commercially available biocompatible coatings, such as non-functionalized parylene, lack

anchor groups, which would allow them to become effective carriers for biomolecules

[20]. Lahann and coworkers [33] prepared a whole range of poly(p-xylylene) films that

have functional groups, which would allow further surface modification. Functional

groups ranged from ketones to carboxyl and amine groups. Thus, a parylene coating

could be functionalized to immobilize a variety of biomolecules, such as sugars and

proteins.

Their excellent electrical insulating properties make organosilicon films useful as

interconnect dielectric material in semiconductors. Their usual deposition technique,

CVD, is compatible with exiting industrial semiconductor technologies, which also

makes this these types of materials a natural choice. The same properties, along with

proven biocompatibility [5, 6, 8] make these types of films good candidates for insulting

coatings in biodevices. Moreover, CVD technique usually employed to obtain siloxane

films enables one-step depositions of the insulating layer along with a layer for

biofunctionalization

17

2. Critical Literature Review

As previously mentioned, reducing dielectric constant values via reducing film

density and polarizability is the goal in semiconductor research at the moment. While

spin-on film preparation allows lower dielectric constant values, mechanical properties of

these films do not always meet industry requirements for film hardness [15] and a high

degree of cross-linking makes them prone to cracks [11]. Also, spin-on film preparation

requires the integration of a wet step into a mainly vacuum process, which can be

impractical. Additionally, spin-on processes are long and require large amounts of

solvents. On the other hand, there are many different ways of obtaining porous thin films

via CVD: incorporation of labile porogen species into the film matrix, deposition of films

using cyclic precursors, and molecular architecture taking advantage of reactive

chemistries in the precursor molecules. Each of the methods has both advantages and

shortcomings. The sacrificial porogen technique is easiest for incorporation of large void

spaces into the film, but it also yields films with poor mechanical and thermal

performance. Intrinsic porosity incorporated into the films as a result of using cyclic

precursors can be insufficient, and the success of this technique is directly dependent on

deposition conditions, since the retention of precursor structure is crucial. The same can

be said for molecular architecture technique: it does not result in significant amounts of

porosity in the films and is very dependent on deposition conditions. Since the focus of

this study is on CVD of low dielectric constant organosilicon thin films, it is important to

learn about the methodologies as well as advantages and disadvantages of each of the

18

techniques for obtaining additional porosity, so as to make a proper choice of one or

maybe even a combination of techniques for optimal results.

2.1

The dielectric material used for spin-on application is typically a biphasic mixture

consisting of the bulk matrix phase, which makes up the ultimate structure of the

resultant film, and the porogen phase, which is the thermally degradable pore-forming

material. The porogen is removed during an annealing step after the film is cast. The

annealing also condenses the matrix phase, giving it more cross-linking, and thus,

structural strength.

Spin-on Technique

Typically, spin-on films are known to have better extendibility to lower dielectric

constant values but also poorer mechanical properties than CVD film, making them less

desirable for industrial applications [35-37]. While increasing cross-linking density

typically improves film hardness, in spin-on films high cross-linking density was also

suggested to lead to cracking during spin coating or annealing when the films reached

some critical thickness [11].

Spin-on films are prepared using SSQ materials [11, 38-40] as well as fully

organic matrix precursor materials [35]. The more traditional SSQ material approach has

been shown capable of yielding films with dielectric constants as low as 2.1 [11, 19].

However this advance in the electrical properties has generally been accompanied by

rather poor mechanical performance: i.e., the low dielectric constant (k=2.1) film

deposited by Kim and coworkers [38] was reported to have a hardness value of only 0.33

19

GPa, while a film with the same dielectric constant deposited by Xi et al. [11] had only a

slightly better hardness of 0.64 GPa.

Research by Ting and coworkers [39] showed the importance of film porosity and

pore size with respect to the physical properties of the film, such as refractive index,

dielectric constant, leakage current, hardness, and elastic modulus, while Takada et al.

[40] related these properties to the skeletal structure of the films as well. One finding in

particular that was of significance concerned the effect that pore generating material had

on pore structure, and, consequently, on film properties [39]. Nonionic surfactant-type

porogen materials, e.g. Brij-56 (polyethylene glycol dexadecyl ether), Tween 80 ((x)-

sorbitan mono-9-octadecenoate poly(oxy-1,2-ethanediyl)), and Pluronic P123

(poly(ethylene oxide)-b-poly(propylene oxide)-b-poly(ethylene oxide)), were shown

to be ideal for obtaining films with optimal mechanical properties, as they yielded

smaller, better organized pore structure within the films. The well-ordered porosity

(porosity of 0.34-0.58, pore size of 2-4.2 nm) within the films also contributed to

improving their mechanical properties. The dielectric constant values ranged between

2.3 and 3.3 and hardness values between 1.22 and 2.52 GPa, as can be seen in Figure 2

and Table 1, respectively. The general trend in all these films was the higher the

dielectric value, the better the mechanical properties. Mechanical properties deteriorated

fast with decreasing dielectric constant values: the dielectric constant decreased from 3.2

to 2.3 when film porosity increased from 45 to 58 % (films prepared using the same

porogen), while the hardness dropped by more than half: from 2.52 to 1.22 GPa. This

observation suggested that there may exist a threshold porosity above which mechanical

20

properties degraded significantly, with only limited improvements in electrical properties

of the film.

Table 1: Hardness (H) and Young’s modulus (E) of each porous film prepared by Ting et al. [39] using different templates

Figure 2: Correlation between dielectric constants and porosities of the films prepared by Ting et al. [39] using different templates.

A singular, innovative approach of preparing spin-on films from organic polymers

was reported by Huang and Economy [35], who were able to obtain films with dielectric

constants as low as 1.9 without compromising the film’s hardness, which was about 2.0

21

GPa. Again, increasing film porosity affected mechanical properties more significantly

than the electrical properties, however, the organic matrix/porogen system seemed to

perform better than the traditional SSQ materials. Though the electrical properties of the

films presented in this paper were the best reported in literature for spin-on films, there

was not enough research done on purely organic systems and their thermal performance

is still suspect.

The common theme in all of the research done in the area of spin-on dielectric

thin films has been the coupling of the electrical and mechanical properties of films:

increasing porosity and decreasing density has been shown to lead to decreasing

dielectric constants, and also to weakened mechanical structures. It appeared that

mechanical structures deteriorated with increasing porosity at a faster rate than the

electrical properties improved. All the research done in the area of spin-on films has

suggested the existence of some optimal porosity range at which a balance can be

achieved between mechanical and electrical properties.

Spin-on films appear to be inferior to CVD films with the same dielectric

constants in terms of their hardness and elastic modulus. Poor mechanical properties

were observed by Hijioka et al. [36] for nonporous spin-on films compared to their CVD

counterparts with nearly the same dielectric values (k=2.8 in case of the spin on film,

contrasted with k=2.9 in case of the CVD film), as seen in Figure 3. The difference in

hardness values was almost threefold: 0.5 GPa for the spin-on film as opposed to 1.5 GPa

for the CVD film. So while spin-on deposition technique allows the ability to obtain

films with low dielectric constants, the mechanical properties of the films are unequal to

those typically exhibited by CVD films. This can be a major issue from the standpoint of

22

integration of these films into production, where they would have to undergo many

mechanically rigorous steps.

Figure 3: Correlation between the mechanical properties: (a) Young’s modulus (E) and (b) hardness (H) and dielectric constant of the low dielectric constant films evaluated by Hijioka et al. [36] using different templates.

Some of the other shortcomings of the spin-on technique include the necessity to

have a wet step in the middle of a vacuum process of chip manufacturing, greater

environmental impact due to the vast use of potentially harmful chemicals and solvents

(e.g. ethanol, methanol, acetone, pyridine [35], methyl isobutyl ketone [38], hydrochloric

acid [35, 39]).

Additionally, no investigators studied the applicability of spin-on films to

multilayer devices: types of devices investigated typically include one layer of dielectric

material and one or two metallization layers. However, this is not the case in industrial

applications, which require circuits to have multiple layers. This means that the spin-on

wet step and annealing step as well as barrier film deposition, metallization, chemical-

mechanical polishing, and other industrial steps need to be repeated several times. This

23

raises the issue of film stability both during the wet step and the high temperature steps.

As the film chemistry needs to be soluble in order for spin-on deposition to work, it may

mean that even after annealing and metallization the film can dissolve or swell in the

solvent during subsequent wet steps. This is a major issue that has not been addressed,

and one that makes spin-on technology rather unappealing from an industrial perspective.

2.2

Sacrificial porogen induced porosity technique for the synthesis of low dielectric

constant films by PECVD is similar in its concept to spin-on methodology. Both

techniques involve the deposition of matrix precursor/porogen mixture and further

annealing of the film in order to remove the labile porogen species leaving void space in

their place, thus decreasing film density and therefore, its dielectric constant. The CVD

sacrificial porogen technique is more evolutionary for the existing technologies, being a

vacuum method, and it eliminates the need for costly and potentially hazardous solvents.

One of the important parameters to consider when employing this technique is thermal

stability of the film matrix. This aspect is often overlooked in spin-on technology as the

wet polymerization steps are always followed by a thermal annealing step, which is

assumed make the film thermally stable to subsequent high temperature steps. Since both

precursor and porogen have to be deposited simultaneously plasma conditions may not be

optimal for both, either resulting in a weakened film matrix or altering the chemical

structure of the porogen molecules, which may affect their lability. If the case is such

that the matrix structure is not cross-linked well enough during deposition, annealing the

film may result in significant thickness loss making the film unusable in industrial

Sacrificial Porogen Technique

24

settings. While it is possible to attain films with low dielectric constant values using this

method, typically the films with lowest achievable dielectric constants also exhibit most

thickness loss as well as worst mechanical properties, which is similar to the trends

exhibited by spin-on films [11, 38, 39].

As it was mentioned before, non-porous CVD films have better mechanical

qualities than spin-on films with the same dielectric constant [36]. This was apparent

comparing the results obtained by Grill [4] and Kim et al. [38]: non-porous films

deposited by each had dielectric constant of about 2.8, however Grill’s film, deposited

from tetramethylcyclotetrasiloxane (H4D4) had a hardness of 1.7 GPa and elastic

modulus value of 16.2 GPa, while the hardness of Kim and coworkers’ SSQ copolymer

bulk film was 1.2 GPa and the elastic modulus: 8 GPa. However, the addition of porogen

to the bulk film deposited by Grill resulted in a drop of not only the dielectric constant

but in hardness and modulus as well: the film with 0.5 porogen ratio (ratio of porogen to

H4D4 precursor flowrate ratio in the gas feed) had a dielectric constant of 2.05 and

hardness of only 0.21 GPa [4]. Thus, it could be seen that the trend in the coupling of

electrical and mechanical properties of the films was the same in CVD as in spin-on

films. In addition to mechanical properties of the films, thermal properties, also observed

by Grill, declined with increasing porogen to matrix ratio: while the thickness of the non-

porous films did not change significantly post-anneal, the film with a dielectric constant

of 2.05 deposited with 0.5 ratio of porogen to precursor in the feed, suffered a thickness

loss of ~25 %, as can be seen in Figure 4.

25

Figure 4: (a) Dielectric constant and (b) thickness change during annealing as a function of the organic porogen fraction in the gas feed during deposition of low dielectric constant films by Grill [4].

The sacrificial porogen CVD method allows the deposition of films with post-

anneal dielectric constants as low as 2.05-2.3 [4, 41-44]. Choice of porogen material,

deposition conditions and annealing temperatures have been shown to influence the

properties of the final films. Both Grill [4], who used two different unnamed proprietary

porogens, and Favennec et al. [43], whose research included hexene, hexadiene, and

cyclohexene oxide porogens, showed the importance of porogen choice. In Grill’s

research, one of the porogens used resulted in dielectric constants only as low as 2.4 (at

porogen fraction of 0.06, with about 30 % thickness loss upon annealing), as seen in

Figure 4(a). Any further increase of porogen fraction resulted in an increase of the

dielectric constant value and even greater thickness loss, due to densification of the film

matrix upon decomposition of porogen species. By contrast, the other porogen allowed

the deposition of films with post-anneal dielectric constants as low as 2.05 (at porogen

fraction of 0.5). In case of this porogen, film thickness loss reached a plateau of 25 %

when the porogen fraction was 0.3, as the forming pores must have retained their

26

structure post-anneal. The importance of choosing the right porogen was demonstrated,

though very little could be concluded about the mechanism that caused the films to

collapse and densify in one case, and for additional porosity to form in the films in the

other, since nothing was known about porogen chemistry. The investigation done by

Favennec et al. [43] compared the effects of different functionalities in organic porogens

(all with six carbon atoms): one or two double bonds in a linear molecule and a cyclic

molecule with an epoxy group. Porogen fraction in as-deposit films was determined by

the examination of the FTIR spectra of the hybrid films and their comparison to porogen-

only and matrix precursor-only FTIR spectra. While the as-deposit films chosen to be

annealed all had different porogen fractions, they all had dielectric constants of 2.8,

which served as a baseline. All three films had only minor thickness losses, less than 5

%, upon annealing, however, the percentage of porogen that remained incorporated was

different in each case, as were the dielectric constants of the post-anneal films. Since

reported post-anneal porogen percentage was a percentage of all material in the post-

anneal films, rather than a percentage of initial amounts of porogen, it was difficult to

conclude how much of the porogen stayed incorporated into the film matrix (as bulk

matrix material could have also been evaporated upon annealing). However, it was

concluded that cyclohexene oxide was the best choice for porogen material, since the film

prepared using this porogen had a post-anneal dielectric constant of 2.2, as opposed to 2.3

and 2.4 for films that were deposited utilizing hexene and hexadiene, respectively. It

appeared logical that using cyclic porogen with an epoxy group would yield the most

porous post-anneal film with lowest dielectric constant, providing that the porogen

retained its cyclic chemical structure during deposition and was incorporated into the film

27

via the cleavage of the epoxy group, since its labile cyclic structure would allow for

generation of the most void space upon annealing. Thus, Favennec and coworkers [43]

showed the importance of porogen selection when using the sacrificial porogen CVD

technique, and in particular the usefulness of epoxy functionality for structure retention

and pore formation in the films after the thermal annealing step.

Both Gates et al. [42] and Castex et al. [44] used cyclic porogen molecules with

epoxy groups in their work in conjunction with cyclic bulk matrix precursor molecules.

Gates and coworkers [42] investigated the effects of substrate temperature during film

deposition (using H4D4 as a matrix precursor and cyclopentene oxide (CPO) as a

porogen) on final film properties. It was observed that lower substrate temperatures

during deposition yielded greater organic content in the films, as evidenced by increasing

intensity of the organic peak between 3100 and 2800 cm-1 in the FTIR spectra with

decreasing substrate temperature, presented in Figure 5. This caused greater changes

upon annealing: structural changes accompanied by thickness loss and decrease in

dielectric constant and refractive index. Thus, the film deposited at 300 °C had a post-

anneal dielectric constant of 2.4 and had undergone a thickness loss of only 5 %, while

the film deposited at 100 °C had a dielectric constant of 2.05 with an extremely high

thickness loss of 48% after annealing. Henceforth, an observation was made that at

higher substrate temperatures it was possible to get a more stable Si-O network in the

films. However, this nearly defeated the purpose of adding organic porogen to the matrix

precursor, since so little of it was incorporated into the as-deposit film to begin with.

28

Figure 5: FTIR spectra of as-deposited organosiloxane films deposited at substrate temperatures of 100, 200 and 300 ºC by Gates et al. [42]. The intensity of the organic peak at 3100-2800 cm-1 increased with decreasing substrate temperature, indicating greater degree of retention of organic moieties.

It was difficult to evaluate just how effective generation of porosity was upon

annealing for each of these cases since the dielectric constants of the as-deposit films

were not provided. It is safe to assume that there was more precursor and porogen

fractionation in case of high temperature deposition, and that the dielectric constant of

that film was higher than the other two both as-deposit and post-anneal. The results [42]

illustrated opposite trends in electrical and thermal properties of films, and suggested that

an optimal deposition space needs to be found where a film with good electrical

properties can be deposited and still be thermally and mechanically robust.

29

Research done by Burkey and Gleason [41] also showed that an improvement in

electrical properties of films via increasing film porosity had an adverse effect on the

thermal properties of the film. The lowest dielectric constant obtained, 2.3, was obtained

at the expense of a loss in thickness as high as 36 %. For comparison, the film with no

addition of porogen had thickness loss of only 5 % after annealing, suggesting that

thickness loss occurred upon removal of the porogen. When porogen decomposition

kinetics were studied by annealing the films at different temperatures between 100 and

425 °C and tracking the concentration of porogen species in the film by relating it to the

concentration of carbonyl groups which were observed in the FTIR spectra, as seen in

Figure 6(a), it was observed that porogen decomposition took place at temperatures of

300 °C and above, and that it took place simultaneously with oxygen cross-link formation

through condensation reactions. Thus, simultaneous porogen decomposition and matrix

cross-linking may not be desirable as the film would not be rigid enough when the pore

formation begins, which could contribute to thickness losses upon annealing. The ideal

situation would be if the porogen were removed at a temperature higher than that needed

for the condensation reactions. Thus, two alternative methods for improvement were

suggested: to search for an alternative porogen with higher decomposition temperature, or

to take advantage of pMMA decomposition kinetics and anneal at such temperatures at

which condensation reactions could still proceed while porogen decomposition would be

slowed down.

30

Figure 6: Extent of porogen decomposition as a function of annealing temperature tracked by (a) comparison of C=O peak in the FTIR spectra and (b) comparison of the normalized area under the C=O peak in films deposited by Burkey and Gleason [41].

Wu and coworkers [45] avoided the typical issues associated with the sacrificial

porogen technique and obtained low dielectric constant films using a novel modified

CVD approach: polystyrene nanospheres were used as self-assembled thermally labile

porogen, thus disposing with the need to deposit simultaneously two different precursors

that have opposite properties. In essence this method combined a traditional CVD

technique with the use of self-assembled porogen moieties, typical of spin-on

methodologies. And even in this case the mechanical and thermal properties of the film

were unknown leaving room for doubt.

Thus, it is difficult to deposit a dielectric film with good electrical, mechanical

and thermal properties using the sacrificial porogen technique. A porogen/matrix pair

needs to be selected that will have proper interactions during deposition, so as to allow

the porogen to become incorporated into the film while still retaining its labile properties

31

to yield void spaces upon annealing [4, 43]. Also, it is preferable for the porogen to

decompose after cross-linking has occurred to support the forming void spaces in the film

[41]. Some of the deposition variables that can have an effect on film properties are

deposition power [41], reactor pressure [44], and substrate temperature [42]. Powers,

pressures and temperatures that are too high appear to cause greater precursor

fractionation in the plasma, thus yielding denser, better cross-linked films, that don’t have

too much organic content, which means that they do not have many labile groups present,

so no additional porosity would form after the annealing step [41, 42]. It can be difficult

to narrow down this optimal deposition space since in each case the electrical and optical

properties of the films track similarly, while the mechanical and thermal properties have

opposite trends. A film with a very low dielectric constant has very poor thermal stability

and hardness, and vice versa, a thermally and mechanically robust film is plagued with

poor electrical properties. So while this technique is quite useful, since it allows creating

additional porosity in the films, it also has many challenges where optimal deposition

conditions are concerned.

2.3

In order to avoid the need to deposit bulk matrix precursor and porogen species,

which typically require opposite deposition conditions to successfully serve their

purpose, simultaneously, the idea of exploiting the intrinsic porosity of cyclic molecules,

such as siloxane rings with organic side chains, was introduced. The key to this

technique was to find the deposition space that would allow for the preservation of

precursor ring structure.

Intrinsic Porosity Technique

32

Tada and coworkers [16] explored the effects that molecular structure of the

precursor has on physical and chemical properties of plasma-polymerized porous

organosilicon films and studied conditions necessary for the retention of cyclic precursor

structure. To find precursors that would best allow for structure preservation during

deposition, six membered Si-O ring-type precursors with various side-chain substitutions

were tested, and the effects of the side chain chemistry were tracked. It was found that

vinyl side-chain chemistry lead to an increase in the deposition rate of the film, and

caused the ring structure and the alkyl side-chains to be preserved, thus leading to more

intrinsic porosity of the film and lower dielectric constant, both of which trends are

demonstrated in Figure 7. This is logical, since the activation energy of the C=C pi

bonds is lower than that for the C=C or C-C sigma bonds in the precursor, 256 kJ/mol

and 346 kJ/mol respectively [46], so the pi bonds in the vinyl groups would be the first

ones to react. Also, it was found that larger alkyl side-chains increased the deposition

rate and the organic content of the film, which yielded lower dielectric constant values,

also illustrated by Figure 7. These observations [16] suggested that low plasma power

and high precursor concentration aided structure preservation, whereas high plasma

powers caused precursor fractionation and loss of the ring structure. This was concluded

based on the dielectric constant values increasing with increasing plasma powers and

decreasing monomer partial pressures, as can be seen in Figure 8(a) and (b), respectively.

The authors suggested using a ring-type siloxane precursor with both unsaturated

hydrocarbon (vinyl) and large alkyl side-chains under low-power plasma with a high

partial pressure of the precursor for optimal results. The lowest dielectric constant

reported by Tada and coworkers [16] was 2.45 obtained using a Si-O six-membered ring

33

precursor molecule with vinyl and the longer alkyl side-chains, at plasma power of 100

W and precursor partial pressure of 0.48 Torr. This value is significantly lower than the

values obtained for dense CVD organosilicon films: 2.8 for a dense H4D4 film deposited

by Grill [4], and 2.9 for a dense siloxane film studied by Hijioka et al. [36], showing that

it is possible achieve a porous low dielectric film utilizing the intrinsic porosity

technique.

Figure 7: (a) Deposition rate and (b) dielectric constant of films deposited by Tada et al. [16] as a function of the partial pressure of cyclic siloxane monomer. The monomer denoted as 3E3RL had three ethyl and three alkyl side chains. The monomers 3V3RS and 3V3RL had three vinyl and three alkyl side chains, where RL side chain was longer than RS.

These results suggested that the key element in this technique is the ability

to preserve the original structure of the monomers. This could come to be an

issue during thermal treatments of the films, an important aspect not addressed in

this particular study.

34

Figure 8: Dielectric constant of 3V3RL films deposited by Tada et al. [16] as a function of (a) plasma power and (b) monomer partial pressure.

The utility of vinyl chemistry as well as the comparison between cyclic and

acyclic precursor chemistries was performed by Burkey and Gleason [14] who deposited

films from V3D3, H4D4, divinyldimethylsilane (S1), and trivinylmethilsilane (TVMS).

It was observed in the FTIR spectra of these films, presented in Figure 9, that the films

deposited from the cyclic precursor had a structure predominantly made up of tri-

substituted silicon atoms (T groups, peak centered at around 1270 cm-1), the TVMS film

had an equal proportion of di- (D, peak centered at around 1260 cm-1) and tri-substituted

groups, while the S1 film had more D than T groups. This was due to the fact that the

cyclic precursors already have a di-substituted structure, thus, they form T groups more

easily, than the linear silane-based precursors with no oxygen in their structure. The

greater reactivity of vinyl groups was apparent since the TVMS films had significantly

more T groups in their structure than the S1 films, and the only difference in the

precursor structure was a third vinyl group in TVMS molecules.

35

Figure 9: The Si(CH3)x stretching region of the FTIR spectrum for (a) TVMS, (b) S1, (c) V3D3, and (d) H4D4 films deposited by Burkey and Gleason [14]. The films deposited from cyclic precursors exhibited the prevalence of tri-substituted silicon atoms, a more mechanically robust structure compared to di-substituted groups.

This difference in film structure was most apparent in low power films (200 W at

10/90 duty cycle), since at high deposition powers precursor fractionation was

responsible for all films having similar structure.

All films had excellent thickness retention upon annealing, >90 %, both the ones

that were deposited at high and low powers, however, those deposited from cyclic

precursors performed better when deposited at lower powers, with thickness retention

being 96-97 %. The H4D4 films exhibited the best mechanical properties both after low

36

and high power depositions, and were the only ones, the mechanical properties of which

improved after the thermal annealing step. This was due to H4D4 molecules preserved in

the final film being linked together upon annealing via condensation reactions, forming a

cage SSQ-like structure, which increased mechanical strength of the films. The V3D3

ring structure, being smaller and more strained, does not have the appropriate substituents

orientation to allow this to happen, and the silane-based acyclic precursors have no

inherent structure that would result in this type of final film matrix. The low power films

had low dielectric constants but their mechanical properties suffered, only the H4D4 film

had a low dielectric constant of 2.4 and a marginally good hardness of 0.54 GPa.

Amongst the high power films the V3D3 film looked promising with a dielectric constant

of 3.0, hardness of 1.3 GPa, and elastic modulus value of 9.7 GPa. However, the high

dielectric constant of this film suggests precursor fractionation and loss of the cyclic

structure.

Burkey and Gleason [15] also utilized a vinyl and methyl substituted siloxane

precursor with intrinsic porosity, specifically V3D3, to deposit thin dielectric films using

pulsed-plasma chemical vapor deposition. They also relied on the reactivity of vinyl

bonds to retain the ring structure of the precursor: the vinyl groups were to be completely

removed due to the disparity in electron density between the silicon atom and the vinyl

group upon reacting with water molecules, which were included to promote the formation

of Si-OH moieties in as-deposit films that could undergo condensation reactions upon

annealing to form a better cross-linked silicon-oxygen network.

37

Figure 10: FTIR spectra of films deposited from V3D3/H2O at three different duty cycles by Burkey and Gleason [15]. Samples deposited at the highest duty cycle (10-40) showed the greatest degree of –OH incorporation.

It was observed that oxygen incorporation into the films increased with increasing

plasma duty cycle (ratio of plasma on-time/total cycle time), since there was more

precursor fractionation, illustrated by the increase of intensity of the hydroxyl peak,

~3600-3200 cm-1, in FTIR spectra of the films in Figure 10.

The mechanical properties of the films tracked directly with the amount of

incorporated oxygen, since higher oxygen content meant better cross-linking in the films,

and also improved upon annealing, since it lead to better film cross-linking via

condensation reactions taking place between the hydroxyl groups. The post-anneal

thickness retention was >95 % in all cases. Dielectric constant values also reflected the

38

degree of –OH incorporation in the films: they increased with increasing duty cycle for

both post-anneal and as-deposit films. The lowest dielectric constant value reported by

Burkey and Gleason [15] for a V3D3-H2O film in this study was 2.6, significantly lower

than the values obtained for dense CVD organosilicon films, which are about 3 [4, 36].

This study was yet another one that pointed to the tradeoffs between mechanical and

electrical properties of dielectric films.

The films deposited by Ross and Gleason [17] proved the usefulness of vinyl

chemistry, as they were deposited from octamethylcyclotetrasiloxane (D4), which only

had methyl side-chains, using hydrogen peroxide as an oxidant, and had post-anneal

dielectric constants of 2.78-3.2, but mechanical properties were inferior to those of films

deposited by Burkey and Gleason [15]. This most likely was due to the fact that the lack

of vinyl chemistry led to the breakdown of the precursor ring structure.

It has been shown [14-17] that low dielectric constant porous films could be

deposited utilizing precursor molecules with cyclic structure, thus, exploiting their

intrinsic porosity. Dielectric constants as low as 2.4 have been achieved in film deposited

via the intrinsic porosity technique [14, 16], and while the sacrificial porogen technique

yielded films with better electrical properties [4], intrinsic porosity films proved much

more thermally stable [14, 15]. However, in order for the intrinsic porosity of the

precursors to contribute to final film porosity and to help lower the dielectric constant of

the film it is crucial to retain the cyclic structure of the precursors. It was seen that in

order to preserve precursor ring structure several steps can be undertaken: precursors with

vinyl and larger alkyl side-chains should be chosen, low plasma power conditions and

high precursor partial pressures should be used [14, 15, 16]. Also, larger ring structures,

39

such as H4D4, may be preferable to smaller ones, i.e. V3D3, since the bonds are less

strained and thus require more energy for fractionation, making preservation of the cyclic

structure easier [14]. At a given hardness value, films deposited from cyclic precursors

appear to have lower dielectric constants than films deposited from acyclic precursors,

which is due to the greater proportion of T groups in these films. This gives the films

greater mechanical stability, while film density is still low due to retained intrinsic

precursor porosity [14].

2.4

The most important thing in implementing the intrinsic porosity technique is

ability to preserve the precursor structure. Several different methods can be utilized in

order to preserve the cyclic structures, such as employing reactive side-chains [16] or

using bigger cyclic molecules [14]. In their study using the sacrificial porogen technique,

Castex and coworkers [44] observed that the addition of cyclohexene oxide porogen to

decamethylcyclopentasiloxane (D5) plasma resulted in improved cyclic structure

retention in the final films. This was attributed to more efficient fragmentation of the

porogen moieties lowering the energy available for dissociation of the D5 molecules.

Thus, addition of a secondary precursor can help preserve the cyclic structure of the

primary precursor by lowering the energy available for fractionation. Also, the addition

of a secondary linear precursor could help spacing the primary precursor molecules

farther apart, easing the strain on smaller cyclic molecules, such as V3D3, thus making

the energy requirement for fractionation greater [14]. Some of the requirements for a

secondary precursor would be that it is more reactive than the cyclic precursor, in order to

Molecular Architecture Technique

40

shield the structure of the primary matrix precursor, and that it form a mechanically

sound densely cross-linked matrix together with the primary precursor. Good candidates

for this type of secondary precursors are linear organosilicon molecules that form

successful thin films with low dielectric constants on their own.

Vinyl chemistry has been shown to be beneficial in linear precursors as well as

cyclic precursors. Millela and coworkers [47] deposited successful low dielectric

constant films from divinyldimethyldisiloxane (DVDMDSO) with post-anneal dielectric

constants reaching as low as 2.1 for films deposited with no oxygen addition. Usefulness

of the vinyl groups was made apparent when DVDMDSO and hexamethyldisiloxane

(HMDSO) films were deposited at the same deposition conditions with a small flowrate

of oxygen [4-16 sccm]. The film deposited from the precursor with vinyl chemistry

retained more organic content and thus had a lower dielectric constant, 2.3, post-anneal,

than the film deposited using the precursor with saturated organic groups, the dielectric

constant of which was 2.9 post-anneal. The thermal stability of the two films also

tracked with the amount of organic content: the more organic content the less stable the

film. Thus, the DVDMDSO film exhibited 11 % thickness loss upon annealing as

opposed to 7 % that the HMDSO film underwent.

Table 2: Comparison between electrical and thermal properties of films deposited from DVTMDSO and HMDSO by Milella et al. [47]

41

Another study examining the reactivity of vinyl chemistry was done by Heo and

coworkers [48], in which films were deposited from precursors with one, two, and four

vinyl groups: vinyltrimethylsilane (VTMS), divinyldimethylsilane (S1), and

tetramethylsilane (TMS), with addition of oxygen gas. While this study did not include

measurements of electrical, mechanical or thermal properties of the films it inspected

film porosity and density, and observed that a silane precursor with a higher number of

vinyl groups resulted in a denser dielectric film, and that a lower population of pores was

generated in films prepared from precursors with higher number of vinyl groups. Since

more vinyl groups lead to higher cross-linking density in the films, and thus, to better

mechanical and thermal stability, the greater density loss upon annealing was in the films

deposited from precursors with less vinyl groups. Hence, yet again, there is a trade-off:

the films that are better cross-linked and more stable are denser and do not show

improvement of properties post-anneal. That is to say, films that exhibit excellent

mechanical and thermal performance are inferior from the perspective of electrical

properties.

It has been shown that the greater the degree of oxygen substitution on silicon

atoms in the film, the better their mechanical and thermal stability [14], so precursors that

have oxygen in their structure would be a logical choice as secondary precursors in

conjunction with cyclic molecules, to help strengthen the film structure. Casserly and

Gleason [49] explored the effect of oxygen substitution on silicon atoms in precursor

molecules, using trimethylmethoxysilane (TMMOS), dimethyldimethoxysilane

(DMDMOS), and methyltrimethoxysilane (MTMOS) precursors for depositions. The

dielectric constants of the films ranged from 2.78 to 3.2, increasing with increasing

42

number of methoxy groups in the precursor, yet again pointing to the relationship

between electrical properties and film stability. Films were also deposited using the

precursors together with oxygen or hydrogen gas. The results suggested that while both

additional gases caused an increase in the dielectric constant, the addition of hydrogen

caused selective activation of the methoxy bonds and formation of silanol groups in their

places, thus preserving the methyl groups, and yielding only a marginal increase in

dielectric constant values. This difference may have had a significant effect upon the

thermal performance of the films, which unfortunately was not studied.

Some common traits that emerged in the studies listed as well as those that

focused on using cyclic precursors included loss of organic groups and increasing

dielectric constants with increasing oxidant flowrate, substrate temperature and

deposition powers [15, 17, 47, 49]. Dielectric constants were shown to decrease upon

annealing of the films, due to improved cross-linking via condensation reactions [15, 47].

The easily activated vinyl chemistry proved to be useful [47, 48], however it needs to be

employed in moderation in order to achieve a balance between good cross-linking and

organic group retention [48]. Also, a siloxane backbone in the linear precursors

improved film stability, as it depended on the degree of oxygen substitution on the silicon

atoms in the film [15, 49], and having Si-O chains in the precursor would eliminate the

need for oxidant addition, which could contribute to the fractionation of the molecules in

the plasma. The compromise between electrical properties and stability of the films was

again shown [47, 49] to be related to film density and cross-linking. Films with higher

cross-linking density are more stable; however, they generally lack organic composition

and perform similar to the glass-like silicon dioxide films currently employed in the

43

technology. Organic group incorporation and pore formation is crucial in order to

achieve lower dielectric constant films, so a balance must be found between too much

organic makeup, which weakens the films structurally, and too dense cross-linking,

which degenerates the dielectric effect of the films.

2.5

PECVD subjects the growing film to ion bombardment, which can result in

damage to film structure, formation of trapped free radicals and dangling bonds in the

film, as well as increase in cross-link density [26, 50, 51]. Avoiding this ion

bombardment is particularly important when trying to preserve intrinsic porosity in the

film, incorporated using delicate structures, such as V3D3. To avoid plasma excitation

and possible damage to the growing film, hot filament CVD (HFCVD) can be used as an

alternative. In HFCVD films are produced by thermal decomposition of the precursor

gas, which is achieved using a resistively heated filament wire, and the radical species

polymerize on the cooled surface of the substrate [51, 52]. The heated filament serves to

drive the decomposition of either the precursor gas, forming the monomer species, or an

initiator, which, in turn, promotes polymerization. Using low filament temperatures

(<800 °C [50]) enables the deposition of films using precursors containing delicate

functionalities. Initiator molecules typically have a thermally labile bond, such as a

peroxy or azo linkage, thus allowing for depositions to take place at low filament

temperatures, at which the precursors are thermally stable [26]. The resulting

polymerization process produces pathways similar to traditional liquid-phase

polymerization [26, 50-52].

Initiated Chemical Vapor Deposition

44

HFCVD technique was used to deposit hexamethylcyclotrisiloxane (D3) and

octamethylcyclotetrasiloxane (D4) films by Pryce Lewis and coworkers [50]. A tantalum

filament array was used with the wires arranged in a parallel array so as to minimize the

thermal gradient. Since no initiator species were used, the filament temperatures during

deposition ranged from 800 to 1200 °C, and there was no deposition observed at filament

temperatures less than 800 °C. The substrate temperature was kept at 60 °C by back-side

water cooling, which was important, since the higher the substrate temperature, the lower

the sticking coefficient of the precursor species, and thus, the slower the deposition. The

deposition rate of both types of film was shown to have an Arrhenius-type relationship as

a function of filament temperature. Comparing FTIR spectra of two D3 films, one

deposited using the PECVD technique, the other, using HFCVD technique at a filament

temperature of 1000 °C, showed more heterogeneity of bonding environments in the

PECVD film, as can be seen in Figure 11. The siloxane peak in the HFCVD film

spectrum was narrower and better defined, while that in the PECVD film spectrum was

wider with a shoulder of greater intensity at higher wavenumbers. Similarly, the organic

peak between 3100 and 2800 cm-1, occurred singularly at ~2965 cm-1, where the methyl

asymmetric stretching peak occurs, with a small symmetric methyl stretching peak at

~2907 cm-1, suggesting possible structural retention of the precursor in the HFCVD film.

By contrast, the spectrum of the PECVD film had a wide organic peak with the presence

of not only methyl but also methylene groups, pointing to fragmentation of the D3

precursor during the plasma deposition. This is hypothesized to be due to the fact that

HFCVD does not provide the same abundance of energy to the monomers inside the

45

reactor chamber as PECVD does, thus resulting in fewer reaction pathways that are

available for the polymerization.

Figure 11: FTIR spectra of (a) D3 PECVD film deposited under continuous-wave excitation and (b) D3 HFCVD film deposited at 1000 ºC by Pryce Lewis et al. [50].

These D3 and D4 films were deposited using three different filament

temperatures: the spectra of the D3 films showed a doublet siloxane peak at all three

filament temperatures (860, 1000, and 1100 ºC), seen in Figure 12 (L), whereas the

spectrum of the D4 film deposited at 800 °C showed only a singlet siloxane peak, while

the films deposited at higher filament temperatures (900 and 1000 ºC) showed behavior

similar to the D3 films, pictured in Figure 12 (R). This suggested that loss of precursor

cyclic structure took place as the filament temperature was increased. While D4 structure

stayed intact at low filament temperature, the D3 ring structure was easily fragmented,

due to the more strained configuration of the six-membered ring, as opposed to the eight-

membered ring of D4.

46

Figure 12: FTIR spectra of (L) D3 films deposited at (a) 860, (b) 1000, and (c) 1100 ºC and of (R) D4 films deposited at (a) 800, (b) 900, and (c) 1000 ºC by Pryce Lewis et al. [50].

The siloxane peak doublet in the spectrum of the D3 film deposited at 1100 °C

resembled that observed in the FTIR analyses of other organosilicon PECVD films. This

suggested that even milder deposition conditions might be needed for incorporation of

six-membered siloxane rings to provide intrinsic pores in thin films.

Initiated CVD (iCVD) was used by O’Shaughnessy et al. [26] as an even milder

alternative CVD technique to deposit V3D3 thin films for biomaterial coatings. t-Butyl

peroxide was used as an initiator, which allowed depositions to be carried out at filament

temperatures of ~200 °C. The comparison of the FTIR spectra of the HFCVD films

deposited at filament temperatures below 500 °C and of the monomer, presented in

Figure 13, showed an almost exact resemblance in the siloxane peak between 1200 and

950 cm-1: the peak was centered at 1012 cm-1 indicating that the siloxane ring structure

47

present in the monomer was retained. There must have been some minimal ring structure

loss, which would explain slight broadening of the siloxane peak in the HFCVD film

spectrum, however the intensity of the higher wavenumber peaks, which could be

identified as long- and short-chain siloxane peaks, was negligible. The oxygen

substitution peak was centered at 1260 cm-1 in both the monomer and the HFCVD

polymer spectra, also suggesting retention of structure, in which each silicon atom has

two bonds to oxygen atoms and two bonds to organic groups. A peak that exhibited the

most change was that at around 1600 cm-1, assigned to C=C stretching in a vinyl bond:

less than 5 % of the absorption intensity remained in the polymer spectrum.

Additionally, absorptions associated with C-H bonds in unsaturated hydrocarbons in the

3100-2800 cm-1 region were significantly reduced in the polymer spectrum. This

indicated that polymerization of V3D3 took place by reaction of the more easily activated

vinyl groups to form saturated linear carbon chains while the rest of the molecule

structure stayed intact.

48

Figure 13: FTIR spectra of the V3D3 monomer and the resulting iCVD polymer deposited by O’Shaughnessy et al. [26]. The presence of the 1012 cm-1 peak in both spectra indicated retention of Si-O ring moieties in the polymer.

The spectra of films deposited at filament temperatures above 500 °C showed a

decrease in intensity at lower wavenumbers, a broadening and a shift of the siloxane peak

towards higher wavenumbers, indicating a loss of the cyclic structure, and a presence of

linear siloxane chains in the films. This suggested that at filament temperatures above

49

500 °C enough energy is provided for polymerization to occur without the addition of

initiator.

Two deposition regimes were observed from the log scale plot of the deposition

rate versus the inverse filament temperature, given in Figure 14: at filament temperatures

of 200-400 °C the deposition process appeared much more sensitive to increases in

filament temperature, at temperatures above 400 °C mass transport limitation seemed to

dominate. However, it was theorized that since the deposition rate was still only about 18

nm/min in this high temperature region, the mass transport limitation occurred not in the

gas phase, but on the substrate due to the adsorption of the monomer onto the substrate

surface. This showed that in iCVD the polymerization process took place on the

substrate itself, as opposed to growth of polymer chains in the gas phase.

Figure 14: Deposition rate data for polymer growth as a function of filament temperature for V3D3 films deposited by O’Shaughnessy et al. [26].

50

Lau and Gleason [52] took this hypothesis further; they theorized that iCVD

polymerization occurs through three major steps: (1) thermal decomposition of an

initiator in the vapor phase to form primary radicals, (2) diffusion and adsorption of

primary radicals and monomer from the vapor phase onto the surface, and (3)

polymerization of monomer on the surface via initiation, propagation, and termination

events to form a continuous polymer coating, as illustrated by Figure 15.

Figure 15: Reaction mechanism proposed for iCVD polymerization by Lau and Gleason [52]. Initiator (I) is thermally decomposed in the vapor phase by heated filaments. Primary radicals (R) and monomer (M) are adsorbed onto the surface. Polymerization occurs at the surface to form the polymer (P) coating.

A series of alkyl acrylates, CH2=CH(COOR), where R=CnH2n+1 (n = 1-6), were

deposited. The FTIR spectra of the iCVD films matched the spectra of films obtained

from the same monomers via liquid-phase radical polymerization. Deposition conditions

were kept constant for all the monomers, and, since the length of the alkyl chains varied,

the only variable was the saturated vapor pressure of the monomer, Psat. Since Psat was

significantly lower for the precursors with longer alkyl chains, it was hypothesized that

they would be more readily adsorbed on the substrate surface at equal monomer gas

pressures, PM. It was proposed that the ratio PM/Psat could be used as a measure of the

amount of monomer adsorbed onto the substrate surface, similar to Brunauer-Emmett-

51

Teller isotherm equation. Deposition rate and molecular weight of the polymer was

shown to be inversely proportional to 1/Psat, as can be seen in Figure 16. This, in turn,

confirmed that the monomer concentration on the surface was the rate-determining

parameter in iCVD.

Figure 16: Effect of monomer-saturated vapor pressure on (a) polymer deposition rate and (b) number-average molecular weight for iCVD poly(alkyl acrylate) films deposited by Lau and Gleason [52]. Ethyl (2A), n-propyl (3A), n-butyl (4A), n-pentyl (5A), and n-hexyl (6A) acrylate data are shown.

The effect of PM/Psat factor was further explored: ethyl acrylate films were

deposited at different substrate temperatures, since, according to the hypothesis of the

authors, a lower substrate temperature would lead to greater monomer absorption, and

thus, to faster deposition rates. Each set of depositions was carried out at three different

filament temperatures: 285, 310, and 360 ºC. It was shown that at each filament

temperature the deposition rate increased appreciably with a decrease in substrate

temperature, indicating adsorption-limited rather than reaction-limited kinetics.

Additionally, the deposition rate data was regressed to an Arrhenius relationship with a

52

slope that gave an apparent activation energy of -79.4 ± 4.7 kJ/mol, presented in Figure

17, which was the same for all three sets of data obtained at different filament

temperatures, suggesting that the same limiting behavior had occurred regardless of

filament temperature. And while deposition rates also tracked with increasing filament

temperatures, the increase was modest, compared to the increase with decreasing

substrate temperatures.

Figure 17: Effect of substrate temperature on polymer deposition rate of iCVD ethyl acrylate films deposited by Lau and Gleason [52] at filament temperature of () 285, () 310, and () 360 ºC. Plotted at a function of substrate temperature each set of data could be fitted into an Arrhenius form.

Monomer choice and substrate temperature were shown to be crucial deposition

parameters in iCVD. Monomers with low saturation pressures would be more readily

adsorbed and would yield faster deposition rates, the trade-off being the difficulty of

generating enough backing pressure to deliver the monomer into the reactor. Lower

substrate temperatures would enhance monomer adsorption and would also lead to a

higher deposition rate; however, there would be a lower limit to substrate temperature

53

used: the saturation temperature of the monomer, at which it would condense. An

optimal range for PM/Psat ratio was suggested to be 0.4 – 0.7 for the case of ethyl acrylate,

at which there would be 1-3 monolayers of monomer at the surface to provide enough

monomer for a suitably fast deposition rate.

2.6

CVD organosiloxane films are of significant interest in the field of implantable

biodevices since they combine biocompatibility with insulating properties. This is

particularly important in the field of neuroprosthetics where the device has to not only

thrive in a biological in vivo environment but also have good dielectric properties so as to

allow the device to utilize the firing patterns of individual neutrons without any

interference. CVD organosiloxane films have been shown to perform successfully as

biomaterials in the past. Chawla showed that coating microporous polypropylene

membranes with PECVD hexamethylcyclotrisiloxane [5] and

octamethylcyclotetrasiloxane [6] films decreased the number of platelets and leukocytes

adhering to the membrane surface, as well as minimized the morphological changes in

the blood cells, which could lead to platelet aggregation, inflammation and thrombosis.

Hasirci [54] showed an improvement in hemoperfusion issues associated with hemolysis

and blood cell adhesion in activated charcoal substrates after the substrates were plasma

coated with hexamethyldisiloxane (HMDSO). The HMDSO coating resulted in a

decrease in hemolysis, breakage of red blood cells to release hemoglobin, from 0.035 to

0.025, and a decrease in percentage of platelet reduction from 29 to 1.6 %. Thus, the

organosiloxane thin film coating was shown to reduce damage to red blood cells in

Organosilicon Films as Biocompatible Insulators

54

biodevices. Ishikawa et al. [7] showed 20 to 50 % improvement in thromboresistance in

glass slides coated with PECVD organosiloxane films: tetramethylorthosilicate,

methyltrimethoxysilane, tetraethylorthosilicate, and others. This suggested that the

biocompatibility of implantable devices could be improved by PECVD organosiloxane

coatings.

Initiated CVD trivinyltrimethylcyclotrisiloxane (V3D3) films were investigated

by O’Shaughnessy and coworkers [8] as possible candidates for biopassivation coatings

on neuroprosthetic devices. These films were shown to have electrical resistivity that

was comparable with that of the current electrical insulation material of choice, parylene-

C: about 4 x 1015 Ω-cm compared with 1 x 1015, and thus a thinner V3D3 would provide

the same level of electrical insulation to the device. The films were shown to be very

flexible, qualitatively, via a scanning electron micrograph, which was an improvement on

the brittle parylene-C. V3D3 films showed no solubility in either polar or non-polar

solvents due to their excellent cross-linking, which was demonstrated by 30 minute soaks

in various solvents, from deionized water to tetrahydrofuran, with no significant

thickness loss. The stability of the films’ electrical properties was proven by monitoring

the electrical resistance of the films under simulated bioimplanted conditions and

constant electrical bias, which did not change over a period of greater than two and a half

years. Lastly, the compatibility of the films with neurons was assessed by culturing PC12

neurons in the presence of glass slides coated with V3D3, and an observation was made

that contact with poly(V3D3 ) did not affect the growth characteristics of PC12 neurons

due to factors such as cytotoxic chemical groups, entrained monomers, or unreacted

initiator, as illustrated by Figure 18. Overall, a conclusion of was made that poly(V3D3)

55

films were promising biopassivation coating materials due to their excellent electrical

resistance, which proved stable long-term even in bioimplanted environment, their good

adhesion and negligible solubility, and their biocompatibility [8].

Figure 18: PC12 neuron growth in the presence of glass substrates both uncoated and coated with iCVD V3D3 by O’Shaughnessy [8]. No significant difference in cell growth was observed, indicating that V3D3 coating was non-cytotoxic to PC12 neurons.

Plasma polymerization of acrylic acid has been shown to provide highly

functionalized surfaces, with the highest retention of the monomer structure and the

highest density of carboxyl groups [25, 34]. The surface density of functional groups is

of great significance, since the density of immobilized bioactive species significantly

affects the sensitivity, detection limits, as well as signal to noise ratio of the biosensor

[34].

Protein adhesion on PECVD acrylic acid films has been shown by Rossini and

coworkers [22]. It was shown that protein attachment took place at carboxylic acid

functional groups, since the FTIR measurements of NH absorption bands characteristic of

protein presence indicated that adhesion occurred only on the films deposited at low

power of 20 W. The increase of power resulted in a decrease of concentration of COOH

56

groups in the film, as evidenced by X-ray Photoemission Spectroscopy (XPS) and

Fourier Transform Infrared Spectroscopy (FTIR), and an increase in the CO

concentration in the plasma phase, observed by Mass Spectroscopy (MS) and Optical

Emission Spectroscopy (OES). The decrease in the concentration of the carboxyl groups,

in turn, resulted in the decrease of hydrogen bonding, hydrophilicity, and acid-base

character of the film. No protein adhesion was observed on the films deposited at 60 W.

This pointed to the fact that films deposited at higher powers lost too much of their

functionality.

Candan and colleagues [25] have shown that the parameter that controlled the

plasma polymerization of acrylic acid was plasma power/monomer flowrate ratio, under

conditions of constant pressure, electromagnetic field, and reactor geometry. This ratio

defined the electron density, which increased with deposition power and monomer

residence time, which, in turn, was inversely proportional to the monomer flowrate. As

power/flowrate parameter increased, the concentration of carboxyl groups in the films

decreased due to more extensive fragmentation taking place in the plasma, greater fluxes

arriving at surfaces in contact with plasma, and more energy being deposited per ion at

the surface. However, while the low-power films retained more functional groups, they

were insufficiently cross-linked to prevent dissolution, which was demonstrated by

washing in distilled water. Deposits obtained at higher powers were better cross-linked

and thus, not easily dissolved. Film stability in liquids is important, since attachment of

biomolecules to the functionalized surfaces is carried out in liquid media. Thus, there

appeared a balance between a sufficient degree of cross-linking, which would prevent

57

film dissolution in liquids, and the concentration of functional carboxyl groups on the

surface.

Akkahat and Hoven [34] obtained linear and branched poly(acrylic acid) brushes

tethered to the surface of silicon oxide substrates via solution polymerization. In case of

linear poly(acrylic acid) (pAA) brushes, the density of carboxyl groups was shown to

increase as a function of chain length, which depended on the monomer-to-initiator ratio.

For branched pAA brushes the concentration of carboxyl groups was shown to decrease

with increasing cross-linking. The authors showed that there is an inverse correlation

between the concentration of functional carboxyl groups and the degree of cross-linking

in the polymer.

2.7

The electrical, optical, mechanical, and thermal properties of low dielectric

constant organosilicon films all directly depend on the chemical structure of the film.

This fact makes the ability to learn about the structure of films crucial to any analysis.

This can be done via Fourier transform infrared spectroscopy. In addition, a way to learn

about the electrical properties of a film is to know what its optical properties are by

performing a spectroscopic ellipsometry analysis. Ellipsometry also provides film

thickness data, and by comparing as-deposit and post-anneal film thicknesses one is able

to learn about its thermal properties.

Characterization Techniques

58

2.7.1 Fourier Transform Infrared Spectroscopy

Fourier transform infrared spectroscopy (FTIR) is a widely used method for

characterization of the chemical composition and chemical bonds in low dielectric

constant films. The technique is based on the absorbance of infrared radiation by

different chemical bonds at different wavelengths and to different extents [1].

Electromagnetic radiation is characterized by its wavelength. Radiation with

wavelengths between 1mm and ~1µm is referred to as infrared radiation. Wavenumbers,

the units most commonly used in IR spectroscopy, refer to the number of waves per unit

length, typically a centimeter, and can be calculated by:

˜ ν = 1λ

(5)

Where

˜ ν is the wavenumber in cm-1, and λ is the wavelength in cm.

The input of energy from the electromagnetic radiation during IR spectroscopy

can cause vibrational or rotational excitation in molecules of materials being analyzed

depending on the frequency of the light. Molecules can absorb rotational or vibrational

energy, thus certain groups can transition from one energy state to another. Every

material is permeable to electromagnetic radiation over a wide range, while at certain

wavelengths it can be absorbent. So in order for the energy to be absorbed, the frequency

of the incident light must correspond exactly to the energy difference between the two

energy states concerned:

E1 − E2 = hν (6)

Where E1 and E2 are the two energy states, h is the Planck constant, and ν is the

frequency of the light. The frequency can be related to the wavelength of radiation with

the following equation:

59

ν =cλ

(7)

Where c is the speed of light and ν and λ are the frequency and the wavelength of

radiation.

Vibrational excitation of molecules takes place due to the fact that the molecule

transitions from lower energy level to a higher one upon absorbing energy, followed by a

release of this energy, which causes the molecule to transition back into the lower energy

state. When the input energy is lower than what is needed for vibrational excitation of

the molecule rotational excitation occurs [55].

Fourier transform spectrometers utilize Michelson interferometers to quantify

radiation intensity variations due to absorbance by the analyzed material using phase

information. The interferometer consists of a beam-splitter that splits the radiation beam

emitted by the source into two partial beams, which are then reflected on a fixed mirror

and on a movable mirror, and then recombined back onto the beam. Shifting of the

movable mirror can cause phase differences between the two partial beams, resulting in a

change in interference magnitude. The changes in optical path length are picked up by

the detector. For a monochromatic light source a cosine signal is obtained at the detector

as a function of the optical path difference x, also referred to as retardation:

I x( )= I0 1+ cos 2π ˜ ν x( ){ } (8)

Where I is the detected beam intensity, I0 is the partial beam intensity,

˜ ν is the

wavenumber of the light coming from the source, and x is retardation. For sources with

more than one wavelength, the interference pattern is calculated by taking the sum of the

cosine signals of all individual frequencies. The interferograms obtained can be

converted into a spectrum using a Fourier transform [55]:

60

S ˜ ν ( )= I x( )cos 2π ˜ ν x( )dx−∞

∞

∫ (9)

FTIR is utilized to track structural changes in films with changes in deposition

parameters, or following an annealing step [4, 15-17, 41-44, 47, 49, 56]. Beer-Lambert

law could be utilized to link absorbance measured by FTIR to concentration of specific

species [41]:

A= log I0

I= ε C[ ]l (10)

Where A is absorbance, I0 is the intensity of the laser source, I is the intensity of

the light that has passed through the sample, ε is the molar absorptivity of the material,

[C] is the concentration of the species, and l is the beam path length. ε and l would

remain constant for a given material and apparatus, thus absorbance is directly

proportional to the concentration of the species.

61

400900140019002400290034003900

Wavenumbers (cm^-1)

-OH

-CHx

Si-H

OxSi-(CH3)(4-x)

Si-O-Si

Figure 19: An FTIR spectrum typical of organosiloxane films.

There are several significant absorbance bands typical of FTIR spectra of

organosilicon films: the most prominent peak occurs between 1200 and 1000 cm-1 and is

indicative of Si-O-Si chains. This wide band is usually deconvoluted into several,

between two and four, disparate peaks: ~1010-1025 cm-1, ~1040 cm-1, ~1085 cm-1,

~1110-1135 cm-1, and ~1185 cm-1. The peak at the lowest wavenumber, ~1010-1025 cm-

1, is usually said to point to siloxane ring structure in the films [41, 42]. The peak around

1040 cm-1 is generally attributed to long-chain Si-O-Si bonding [41, 49], while the peak

around 1085 cm-1 is said to result from the absorption by long-chain Si-O-Si terminated

with methyl or hydroxyl groups [41, 49]. However, sometimes this area of the spectrum

is viewed as only one peak around 1050-1063 cm-1, which is then assigned to siloxane

networks with Si-O-Si angles ~144° [16, 42]. The peak around 1110-1135 cm-1 is

62

interpreted as short-chain Si-O-Si [41, 49], as cage-type structure seen in spin-on films

[41, 42, 49], or as Si-O-Si bonds with a larger angle, ~150° [41, 42, 56]. And lastly, the

~1185 cm-1 peak is attributed to short siloxane chains terminated by methyl or hydroxyl

groups [41, 49].

The next region of importance in FTIR spectra of organosiloxane films, between

1240 and 1300 cm-1, shown in Figure 20, contains only one peak that signifies the

bonding environment of silicon atoms in the film. There are four possible environments:

mono-, di-, tri-, and quad-substituted, FTIR absorbance of the first three of which would

appear within this region [14, 15, 19, 41, 42, 49]. The degree of substitution refers to the

number of oxygen atoms that each silicon atom is linked to and the rest of the linkages

are to organic groups, typically, methyl. The exact position of the peak in the region can

be used to determine the degree of oxidation of the silicon atoms: a peak centered at

~1250 cm-1 suggests mono-substituted groups, at ~1260 cm-1, di-substituted groups, and

at ~1270 cm-1, tri-substituted groups [15, 56]. The absorbance from quad-substituted

groups does not appear in this region but rather between 1200 and 1000 cm-1, since there

are silicon-organic bonds in these groups.

63

12401245125012551260126512701275128012851290

Wavenumbers (cm^-1)

D-groupsT-groups

Figure 20: FTIR absorbance peak indicative of the degree of oxidation of silicon atoms.

A few other smaller peaks that are typical of organosilicon spectra are the alkene

peak, C=C, at 1600 cm-1 [15, 16, 42], the peak attributed to methacrylate ester, C=O,

occurring at ~1750-1725 cm-1 [41, 42], and the Si-H peak at ~2200 cm-1 [4, 16, 17, 42].

The band of significant intensity occurring between 2800 and 3100 cm-1 is indicative of

organic groups in the film and can also be deconvoluted into separate peaks specific to

methyl and methylene groups [4, 15-17, 41, 49]. The broad peak around 3100-3600 cm-1

is attributed to the incorporation of hydroxyl groups into the film [14, 15, 17, 41].

A summary of all the typical groups present in a siloxane film and their

corresponding peak positions is given in Table 3.

64

Table 3: FTIR peak assignments for organosiloxane films. Wavenumber (cm-1)

Assignment Reference

3100-3600 -OH stretching [15, 41, 49] 2972-2952 Methyl C-H antisymmetric stretch [4, 41] 2882-2862 Methyl C-H symmetric stretch [4, 41] 2960 CH3 antisymmetric stretch [41] 2936-2916 Methylene C-H antisymmetric stretch [41, 52] 2863-2843 Methylene C-H symmetric stretch [41, 52] 2232 Si-H symmetric stretching in H-SiO3 [4, 16, 42] 2178 Si-H symmetric stretching in H-SiO2Si [41] 2165 Si-H symmetric stretching in H-SiOSi [41] 1759-1720 Methacrylate ester C=O stretching [41, 42] ~1600 Alkene C=C stretching [15, 16, 41] 1460 CH3 antisymmetric bending [41, 42] 1440 CH3 symmetric deformation [41] 1410 Si-CH3 antisymmetric stretch [41] 1390 Methyl symmetric bending in Si-CH3 [42] 1360 Methylene bending in Si-CH2-Si [42] ~1270 Si-CH3 symmetric bending in O3SiCH3 [15, 16, 41, 49] ~1260 Si-CH3 symmetric bending in O2Si(CH3)2 [15, 41, 49] ~1250 Si-CH3 symmetric bending in OSi(CH3)3 [15, 42, 49] 1250-1220 Methylene symmetric bending [42] 1210-1160 C-O stretching bands [41] 1200-1000 Si-O-Si stretching [15, 41] ~1190 Short-chain Si-O-Si with terminal CH3/OH [41, 49] ~1135-1110 Short-chain Si-O-Si or cage structure [4, 16, 41, 42, 49] ~1085-1063 Long-chain Si-O-Si with terminal CH3/OH

Network Si-O-Si angle ~144° [41, 42]

~1050-1035 Long-chain Si-O-Si [16, 41, 49] 1025-1010 Si-O-Si in cyclotrisiloxanes [41, 42] 920-830 Si-OH silicon-hydroxyl stretching [15, 41, 49] 870-750 Si-CH3 silicon-methyl rocking [15, 41] ~800 C-Si-C stretching [41]

65

2.7.2 Spectroscopic Ellipsometry

Spectroscopic ellipsometry is a preferred method in the semiconductor industry

for the measurement of film thickness, refractive index, and extinction coefficient [14,

17, 38, 40, 41, 45, 49, 58]. It is a non-destructive contactless optical technique that

allows for the measurements of sample properties based on the changes in polarization of

the light that is reflected from the sample.

In order to characterize a light wave, or any classical wave, the wave intensity,

frequency, direction of propagation, orientation of vibrations and the variation of all of

these parameters with time need to be known. The orientational characteristics of the

wave in time and space are referred to as the polarization of the wave. “Ordinary” light

has no inherent directional quality nor is it affected by phase delays between orthogonal

components [58]. If all photons of a light beam are oriented in the same direction the

light is referred to as polarized light. This can be achieved using either a source that

emits only polarized (such as a laser) light or a polarizer, which is an optical element that

allows light of only one particular orientation to pass through. Ellipsometry utilizes

elliptically polarized waves: two linearly polarized waves with the same frequency

combined out of phase by 90°. The amount of the elipticity of the light changes when it

is reflected off the sample being studied [59].

When a light beam hits the sample being studied is slows down, changes

direction, and part of it is absorbed. The sample is characterized by its thickness and by

complex index of refraction, Ñ:

˜ N = n − jk (11)

66

Where n is the index of refraction, the relationship between the speed of light in

vacuum versus the speed of light in the material studied, k is the extinction coefficient,

measure of how quickly the intensity decreases as the light passes through the material,

and j is the imaginary unit.

The equations for index of refraction and extinction coefficient are as follows:

n =cv

(12)

k =λ

4πα (13)

Where c is the speed of light in vacuum, v is the speed of light in the sample, λ is

the wavelength of the light, and α is the absorption coefficient.

The value of the refractive index of a semiconductor as a function of wavelength

of the light in the visible region is typically described by the Cauchy function:

n λ( )= n0 +n1

λ2 +n2

λ4 (14)

Where n0, n1, and n2 are called Cauchy parameters. The extinction coefficient can

be modeled by an equation of the same form:

k λ( )= k0 +k1

λ2 +k2

λ4 (15)

Where k0, k1, and k2 are called Cauchy extinction coefficients. Another way of

modeling the extinction coefficient of a film is using the Urbach equation:

( ) ( )bEECeCk −= 21λ (16)

Where

E =12400

λ and

E b=12400

λ0

, λ0 is typically set at the lowest wavelength

measurable, and C1 and λ0 are correlated [59].

67

Ellipsometric measurements may be performed as a function of wavelength, of

the incidence angle, or as a function of both. Generally ellipsometers contain a light

source, a polarizer, to fix the state of polarization of the incident film prior it reaching the

sample, and a detector to measure the state of polarization of the light beam after

reflection. Spectroscopic ellipsometers employ light sources that provide a beam with a

range of wavelengths, as opposed to single wavelength ellipsometers, that typically have

a laser generating the beam [59].

Thus, taking into consideration the findings of many previous investigations, it

was decided to proceed with experiments centered around the molecular architecture

technique, as it appeared to yield films with better mechanical and thermal properties

than the sacrificial porogen technique [14, 15]. Films would be deposited from cyclic

precursors, and porosity would be further enhanced by the addition of linear co-

monomers, which would help cyclic structure retention, as well as add additional porosity

by spacing cyclic molecules farther apart [14, 44]. FTIR spectroscopy and spectroscopic

ellipsometry would be used as primary methods of thin film characterization.

68

3. Experimental Approach

Although spin-on films are capable of having lower dielectric constants, this

technique is still not the most desirable in an industrial setting, since spin-on films have

inferior mechanical properties, and have been shown to be prone to cracking when they

reached a certain thickness. The spin-on coatings are performed as a wet-step while all

the other production steps take place under vacuum. It takes longer than CVD steps, and

has been shown to be more expensive and harmful to the environment due to the large

amount of solvents and chemicals needed. In addition, another shortcoming of this

technique is the need to repeat the process several times to achieve multiple dielectric

layers in the circuit: the same solvent is used during spin coating of new layers which can

cause damage to the already annealed dielectric layers. Therefore, in order to achieve

low enough dielectric constants (< 2.4) in CVD thin films to keep pace with industry

needs, nanoscale porosity must be incorporated directly into organosilicon thin films.

Two main techniques can be used: one, which employs the intrinsic molecular-size

porosity of precursor molecules, and another that induces film porosity through addition

of porogen molecules.

While addition of sacrificial thermally labile porogen molecules typically yields

films with dielectric constants lower than many other techniques, this method has been

shown to be plagued by many issues, such as excessive thickness loss and deterioration

of mechanical properties of the films. This is due to the fact that pores forming due to

decomposition of porogen species upon annealing can collapse if there is not enough

cross-linking support. Porogen molecules can be only partially removed if they become

69

permanently bonded to the matrix, harming the electrical properties of the film, if the

dielectric constant of the porogen species is higher than that of the matrix precursor.

Also, it is possible that if the partly incorporated porogen remained labile and is removed

post-anneal it can cause structural weakening of the matrix at the point where the matrix-

porogen bond is broken. The most critical issue associated with the sacrificial porogen

technique is that of deposition conditions. These conditions, if carefully controlled, that

lead to formation of a robust film matrix (e.g. high plasma power or high substrate

temperature), but can also cause fractionation and loss of labile structure of porogen

species. However, at milder deposition conditions, which allow the retention of the

porogen structure in the films can simultaneously yield films that are not cross-linked

well enough and end up being thermally unstable and mechanically poor.

Figure 21: Incorporation of thermally labile porogen in a bulk matrix.

A way to improve the mechanical properties of the film matrix is to use

precursors with inherent molecular-size porosity, such as V3D3 (Figure 22). The

emphasis in this type of depositions has to be on structure retention of the precursor, thus,

making the vinyl side chains important, as they have been shown to help protect other

sites in a molecule from reacting [14, 16, 47, 56]. Also, it has been suggested that larger

ring structures are more likely to stay intact during plasma depositions [14, 44], as there

70

is less inherent strain in the structure. In terms of deposition conditions: low power,

pulsed plasma, low substrate temperature, high reactor pressure, and high precursor

flowrate have all been said to help the retention of the cyclic precursor structure.

Figure 22: Cyclosiloxane precursors with vinyl and methyl side chain chemistry.

Additional free volume could be incorporated into the thin films by utilizing

linear siloxane molecules as spacers between units of the cyclic precursor. This route

could be made easier by taking advantage of side chain chemistries of the cyclic

precursors, as well as by choosing linear precursors with reactive end groups as well, thus

copolymerizing the cyclic and acyclic precursors together. This technique would allow

better control over film density, which could be achieved by choosing spacer molecules

of different lengths (of siloxane chains), such as divinyldimethylsilane (S1),

71

divinyltetramethyldisiloxane (S3), or divinylhexamethyltrisiloxane (S5) (Figure 23).

Other choices for spacer molecules are methyldiethoxysilane (DEMS) and

vinylmethyldiethoxysilane (VDEMS), due to their varied reactive substitutions on the

silicon atom.

Figure 23: Linear spacer molecules with different functional groups.

72

Figure 24: Schematic representation of organosilicon films with molecular (due to V3D3) and tunable (due to spacer molecules) porosity; 1-atom spacer elements on the left, 5 atom spacer elements on the right.

VDEMS molecules in particular would be interesting to investigate, as the easily

activated vinyl group would be first to react, while the ethoxy groups would allow for

condensation reactions to take place, promoting better cross-linking. These types of films

would utilize both intrinsic porosity and molecular architecture techniques, and would

allow for better film cross-linking without the application of too harsh deposition

conditions. It was important to find and optimize deposition conditions for a precursor or

a precursor-spacer couple so that lowest possible dielectric constant value could be

achieved without the addition of porogen.

All depositions were performed in a custom-built cylindrical vacuum chamber

with a volume of 17.2 L. The reactants were fed into the reactor through the showerhead,

which was suspended ~3.5 cm above the substrate for PECVD depositions, and ~10 cm

for iCVD depositions. The V3D3, VDEMS, S1, S3, and S5 monomers were obtained

from Gelest, tert-butyl peroxyde (tBPO) initiator and acrylic acid reactant were obtained

73

from Aldrich, and ethylene glycol diacrylate (EGDA) was obtained from Monomer-

Polymer and Dajac Labs: all reactants were utilized without further purification. The

monomers were heated in crucibles of either aluminum or Pyrex glass in order to provide

sufficient backing pressure for vapor delivery into the reactor. V3D3 was heated to 70 ±

5 °C. Small amounts of copper (II) chloride or hydroquinone (Aldrich) radical scavenger

were added to the crucible with V3D3 to prevent autopolymerization at the high

temperature. VDEMS was heated to 40 ± 5 °C. The temperatures that S1, S3 and S5

precursors were heated to were 30 ± 5, 25 ± 5, and 80 ± 5 °C. tBPO initiator did not need

to be heated, as it was volatile enough at room temperature. The crucible with acrylic

acid (AA) was heated to 60 ± 5 °C, while the temperature of EGDA cross-linker was

varied from 25 to 75 °C in order to supply different amounts of reactant into the chamber.

Ferric chloride was also added to EGDA to prevent autopolymerization at higher

temperatures. A list of all chemicals used is presented in Table 4.

Figure 25: Molecules used in deposition of biofunctionalizable coating.

74

Table 4: List of chemicals used Chemical name Abbre-

viation Supplier Purity (%) Use

Trivinyltrimethyl- cyclotrisiloxane

V3D3 Gelest 99.1 Cyclic matrix precursor for films with intrinsic porosity

Vinylmethyl- diethoxysilane

VDEMS Gelest >95 Linear siloxane precursor deposited on its own and as a spacer co-polymerized with V3D3 for tunable porosity

Divinyldimethyl- silane

S1 Gelest >95 Linear spacer co-polymerized with V3D3

Divinyltetramethyl- disiloxane

S3 Gelest >95 Linear spacer co-polymerized with V3D3

Divinylhexamethyl- trisiloxane

S5 Gelest >95 <5% of S3

Linear spacer co-polymerized with V3D3

Acrylic acid AA Aldrich 99 Coating for further biofunctionalization

Ethylene glycol diacrylate

EGDA Monomer-Polymer and Dajak Labs

Co-polymerized with AA to act as a cross-linker

tert-Butyl peroxide tBPO Aldrich 98 Initiator

Copper chloride Aldrich 97 Radical scavenger

Hydroquinone Aldrich Radical scavenger

Argon Ar Med Tech

Inert diluent gas

The flowrate of the precursors was controlled by a regular needle valve (Model

SS-4BMG-VCR, Swagelok) that was previously calibrated. The flowrate of Ar gas was

controlled by a mass flow controller (Model 1178A12CR1BV--S, MKS). The pressure in

the reactor was maintained by a butterfly valve (Model 253B-1-40-2, MKS). All films

were deposited on 100 mm diameter p-type silicon wafers with <100> orientation

(Montco Silicon). The substrate temperature was controlled by contact with a stage with

backside water cooling using a recirculating chiller (Model 1175MD, VWR).

75

Continuous plasma power for the PECVD film depositions was provided by the Comdel

CPS-500AS 13.56 MHz RF power source. For iCVD depositions 80/20 Nichrome wire

filaments of 26 gauge (0.40 mm diameter, Omega Engineering) were used in a parallel

array spaced 2 cm apart, designed to minimize temperature gradients between individual

wires and provide uniform heating over the substrate. Filament temperature was

measured to ±50 °C certainty by a thermocouple (type K, Omega Engineering) attached

directly to one of the filament wires. The springs on the filament holder maintained wire

tension to compensate for thermal expansion of the wires and to prevent drooping and

possible shorting out. The filament to substrate distance was kept at 2.5 cm. The

filament array setup is shown in Figure 26.

Figure 26: Photograph of the filament holder used in iCVD depositions.

76

FTIR was used to elucidate the structure and composition of the films. Perkin

Elmer Spectrum 2000 spectrometer was used in normal transmission mode to acquire the

spectra at 1 cm-1 resolution over the range of 4000-400 cm-1. Measurements were

averaged over 32 scans. Since all samples were deposited on silicon substrates, a

background of bare silicon was taken prior to sample acquisition. All samples were

baseline corrected. Since FTIR is a line of sight measurement, the area under a peak

depends on the thickness of the sample. Consequently, the spectra were thickness

normalized to allow for accurate comparison, as the films differed in thickness. The

concentration of specific chemical bonds in the film is proportional to the area under their

characteristic absorbance peak. Thus, it is possible to determine the amount of material

using transmission FTIR. Therefore, FTIR bands were deconvoluted into their

component peaks, to better understand the molecular structure of the films, and the peak

areas were determined by integrating the FTIR bands in the spectra. The deconvolution

was performed using the Solver program in MS Excel, where the FTIR spectra files were

exported. Siloxane bands were split up into 5 peaks: cyclic peak at ~1015 cm-1, long-

chain peak at ~1035 cm-1, long-chain with end-capping –CH3/OH at ~1075 cm-1, short-

chain peak at ~1135 cm-1, and a peak characteristic of short-chains with end-capping –

CH3/OH at ~1185 cm-1. Approximate intensities of each peak in the band were entered

into Solver, after which it computed the exact position and intensity of each peak, so as to

minimize the error between the model and the actual spectrum.

Sample thickness and index of refraction was measured using a variable-angle

spectroscopic ellipsometer: JA Woolam ESM-300. A Cauchy-Urbach model was utilized

to obtain a nonlinear least-squares fit of the data obtained at three angles, 65°, 70°, and

77

75°, and 225 wavelengths. Thermal stability analysis was be done by measuring film

thickness before and after annealing at 400 °C under nitrogen atmosphere. Post-anneal

FTIR spectra were useful in determination if any structural changes occurred in the films

upon annealing.

78

4. Results and Discussion

Depositions were performed using only V3D3 as a monomer, as well as using

silane or siloxane co-monomer in order to prevent V3D3 precursor fragmentation in

plasma. These combinations were utilized to test the hypothesis of ability to tailor film

porosity by varying spacer chain length. Due to the strained cyclic configuration of the

V3D3 molecule low energy CVD methods were utilized, such as PECVD with minimal

plasma powers, an improvement on which was iCVD with minimal hot wire

temperatures. Additionally, the V3D3 films were deposited in combination with AA as

possible biocompatible insulating coatings.

4.1 Plasma Enhanced CVD

While addition of sacrificial thermally labile porogen molecules typically yields

films with dielectric constants lower than other techniques allow, this method is also

plagued by many issues, such as excessive thickness loss and deterioration of mechanical

properties of the films. This is due to the fact that pores forming due to decomposition of

porogen species upon annealing can collapse if there is not enough cross-linking support.

Also, porogen molecules can be removed only partially if they become permanently

bonded to the matrix, harming the electrical properties of the film, if the dielectric

constant of the porogen species is higher than that of the matrix precursor. It is also

possible that if the partly incorporated porogen remains labile and is removed post-

anneal it can cause structural weakening of the matrix at the point where the matrix-

porogen bond was broken.

79

A way to improve the mechanical properties of the film matrix would be to use

precursors with inherent molecular-size porosity. The emphasis in this type of

depositions has to be on structure retention of the precursor, thus, making the vinyl side

chains important, as they have been shown to help protect other sites in a molecule from

reacting [26]. In terms of deposition conditions: low power, pulsed plasma, low substrate

temperature, high reactor pressure, and high precursor flowrate have all been said to help

the retention of the cyclic precursor structure [14-16]. It was hypothesized that additional

free volume could be incorporated into the thin films by utilizing linear siloxane

molecules as spacers between units of the cyclic precursor. This route could be made

easier by choosing both cyclic and linear precursors with reactive side-chain chemistry,

such as vinyl groups, which would then help protect the precursor structure [16, 47]. It

was hypothesized that film density could be altered by choosing spacer molecules of

different lengths (of siloxane chains). Linear precursors that have various reactive side-

chains, such as vinylmethyldiethoxysilane (VDEMS) molecules, in particular would be

interesting to investigate, as the easily activated vinyl group would be first to react, while

the ethoxy groups would allow for condensation reactions to take place, promoting better

cross-linking. These types of films would utilize both intrinsic porosity and molecular

architecture techniques, and would allow for better film cross-linking without the

application of too harsh deposition conditions.

4.1.1 Vinylmethyldiethoxysilane Depositions

The deposition space for each precursor, VDEMS and V3D3, was investigated

and conditions were optimized so that the lowest possible dielectric constant/refractive

80

index value can be achieved prior to depositing combination films. Deposition

conditions for VDEMS precursor were investigated. Since it was determined that plasma

power, substrate temperature and precursor flowrate all had significant effects on the

properties of the final films a three variable two level orthogonal design of experiment

matrix was developed using plasma powers of 20 and 100 W, substrate temperatures of

25 and 200 °C, and VDEMS flowrates of 1 and 5 sccm. The design of experiment matrix

of conditions is given in Table 5.

Table 5: Orthogonal design of experiment matrix Plasma Power (W)

Flowrate (sccm) Stage Temperature (ºC)

20 5 25 100 1 25 100 5 200 20 1 200

All films were deposited at a reactor pressure 250 mTorr and using 10 sccm of

argon as a carrying gas. Ar gas was necessary to facilitate plasma ignition for depositions

done using low precursor flow rates. It was observed that film deposition rate increased

with increasing precursor flowrate and plasma power, and decreased with increasing

substrate temperature.

81

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F061111-A - 20W, 5sccm, 25C

F061111-B - 100W, 1sccm, 25C

F061111-C - 100W, 5sccm, 200C

F061111-D - 20W, 1sccm, 200C

Figure 27: Overall FTIR spectra of the VDEMS films deposited at the

conditions specified by the design of the experiment matrix.

From the FTIR spectra of the four VDEMS films presented in Figure 27 it could

be concluded that substrate temperature played the most significant role in determining

the chemical structure of the films. It could be seen in Figure 27 that the Si-O-Si peak in

the 1200-950 cm-1 region is of much greater intensity for the films deposited at 200 °C.

At the same time these films also had lower organic peak intensities at around 3100-2800

cm-1. This inverse proportionality of siloxane and organic peak intensities suggested that

film cross-linking took place upon the loss of organic moieties in the film. The Si-O-Si

peaks in the high substrate temperature film spectra in Figure 27 were also centered

around lower wavenumbers, indicating the predominance of long siloxane chains [16, 41,

42]. It could be seen in the FTIR spectra in Figure 31 that at high substrate temperature

the film deposited at a lower flowrate and lower plasma power had the Si-O-Si peak of

82

greatest intensity. Comparing the FTIR spectra of the two low substrate temperature

films in Figure 29, the film deposited with low precursor flow rate had a smaller siloxane

peak, but it was more skewed towards lower wavenumbers, indicating more ordered

longer-chain Si-O-Si structure in the films. These observations suggested that there was

more control in depositions done at a low precursor flowrate regardless of the deposition

power. It may be hypothesized that since there is less material in the chamber one would

expect a slower deposition rate, but due to greater ion bombardment of each molecule it

reacted more thoroughly and was better cross-linked into the film matrix.

Better film matrix cross-linking in high substrate temperature films was evidenced

by the FTIR spectrum region attributed to oxygen substitution on silicon atoms, 1280-

1250 cm-1, shown in Figure 28. The film deposited at low temperature, 20 W, and 5

sccm flow rate had majority of tri-substituted groups with only slightly less di-substituted

groups, the other low temperature film had mainly tri- and some di-substituted groups,

while the two high substrate films had much smaller peaks in this region and those that

were present were indicative of tri-substituted groups only. The lower intensity of the

peaks was to be expected since there appeared to be minimal organic content in these

films, meaning that most silicon atoms were most likely quad-substituted with oxygen

atom, and would not show up in the 1280-1250 cm-1 spectral region, but rather in 1200-

950 cm-1, adding to the intensity of the siloxane peak, as was seen in Figure 27.

83

-0.0025

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0.0175

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F061111-A - 20W, 5sccm, 25C

F061111-B - 100W, 1sccm, 25C

F061111-C - 100W, 5sccm, 200C

F061111-D - 20W, 1sccm, 200C

Figure 28: Segment of FTIR spectra of VDEMS films deposited at the

conditions specified by the design of the experiment matrix. The shift of the peak towards greater wavenumbers indicates greater oxygen substitution of the silicon atoms, and thus, a greater degree of cross-linking in the films.

84

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F061111-A - 20W, 5sccm, 25C

F061111-B - 100W, 1sccm, 25C

Figure 29: Segment of FTIR spectra of as-deposit VDEMS films

deposited at low substrate temperature. The breadth of the siloxane peak indicates heterogeneity of Si bonding environments in the film.

The post-anneal spectra of the low substrate temperature films, presented in

Figure 30, looked significantly different from the as-deposit spectra in Figure 29: the

spectrum of the 20 W, 5 sccm film in particular. A shift could be observed in the Si-O-Si

peak towards lower wavenumbers, suggesting formation of longer siloxane chains and

the loss of end-capping methyl groups. The organic peak decreased significantly,

suggesting that upon annealing, as well as during deposition, additional cross-linking

takes place at the expense of organic inclusions.

85

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F061111-A-PA - 20W, 5sccm, 25C

F061111-B-PA - 100W, 1sccm, 25C

Figure 30: Segment of FTIR spectra of post-anneal VDEMS films

deposited at low substrate temperature. The shift of the siloxane peak towards lower wavenumbers indicates formation of longer siloxane chains and cross-linking of the films taking place during annealing.

The spectra of the elevated substrate temperature films, presented in Figures 31

and 32, looked the same as-deposit and post-anneal, implying that no structural changes

took place in those films. This was to be expected as these films had minimal organic

content, so there were not many sites at which further cross-linking could occur. The as-

deposit films were already rigid, glass-like and had a high cross-link density.

The peak in the 1280-1250 cm-1 region also changed significantly for the low

substrate temperature films after the annealing step, seen in the comparison of Figures 29

and 30. The peak decreased in intensity and shifted towards higher wavenumbers with

only a small shoulder at ~1260 cm-1 for the 20 W, 5 sccm film, and while it only changed

86

slightly the 100 W, 1 sccm film spectrum, the change also reflected further cross-linking

having taken place upon annealing. In the spectra for the high substrate temperature

films, Figures 31 and 32 there was little change observed.

Thickness retention of the films post-anneal appeared to directly correlate with

the degree of cross-linking in the as-deposit films, reflected by the degree of oxygen

substitution on silicon atoms: the more tri-substituted groups in the film, the better its

thickness retention.

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F061111-C - 100W, 5sccm, 200CF061111-D - 20W, 1sccm, 200C

Figure 31: Segment of FTIR spectra of as-deposit VDEMS films

deposited at high substrate temperature. The position of siloxane peak at low wavenumbers indicates predominance of long-chain siloxane structure in the film.

87

Table 6: Optical properties and thermal stability data for the VDEMS films deposited at the conditions specified by the design of the experiment matrix

Plasma Power (W)

Flowrate (sccm)

Stage (°C)

Refractive index (AD)

Refractive index (PA)

Thickness Loss (%)

20 5 25 1.416±0.071 1.453±0.002 16 100 1 25 1.470 1.411 7 100 5 200 1.430 1.429 0 20 1 200 1.435±0.002 1.435±0.001 0

-0.01

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bsor

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e (a

rbitr

ary

units

)

F061111-C-PA - 100W, 5sccm, 200C

F061111-D-PA - 20W, 1sccm, 200C

Figure 32: Segment of FTIR spectra of post-anneal VDEMS films

deposited at high substrate temperature.

The effect of annealing on the optical properties and the thermal stability data is

presented in Table 6, which shows that the high substrate temperature films exhibited

neither thickness loss nor a change in the refractive index value upon annealing, which

was consistent with the fact that the FTIR spectra of these films showed no change in the

88

film structure after annealing. While the refractive index values of the low substrate

temperature films were higher than those of the high substrate temperature films as-

deposit, the film deposited using low VDEMS flowrate (1 sccm) and high plasma power

(100 W) exhibited a decrease in the refractive index post-anneal. This is likely due to the

reaction of dangling bonds, which increase the polarizability of the film, further cross-

linking, and the formation of a more ordered film structure that took place during the high

temperature annealing step. It is hypothesized that while some organic content was lost

from the film structure (as suggested by the decrease in intensity of the organic band at

3100-2800 cm-1 in Figure 33(b)) the loss was minimal and did not result in a major

densification of the film, also evident from only a small percentage of thickness lost,

shown in Table 6.

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As-depositPost-anneal

Figure 33: (a) Siloxane and (b) organic ranges of FTIR spectra of as-

deposit and post-anneal VDEMS film deposited using 1 sccm VDEMS flowrate, 100 W plasma power, and 25 °C substrate temperature. Both segments of spectra show only small differences between the films before and after annealing.

89

The other low substrate temperature film that was deposited at low plasma power

(20 W) using a high precursor flowrate (5 sccm) showed an increase in the refractive

index upon annealing. This could be attributed to film densification that occurred during

annealing as organic content was lost and new cross-links formed, which is suggested by

the FTIR spectra of the film before and after annealing presented in Figure 34. The

siloxane band in Figure 34(a) shows a significant shift from high wavenumbers,

characteristic of short Si-O-Si chains, or cage structures, to low wavenumbers,

characteristic of long-chain structure. The loss of cage structure, with its built-in void-

space could be indicative of densification. As could the significant loss of organic

groups, evident from the decrease of the organic band intensity, shown in Figure 34(b), as

terminal methyl groups form void spaces around themselves. The significant loss of

organic groups and densification of the film matrix also explains the loss in film

thickness observed after the annealing step, given in Table 6.

High plasma power (100 W) was shown to yield fairly well cross-linked films,

since even the film deposited at low substrate temperature (25 °C) had only a minimal

thickness loss (Table 6) and minimal changes in its molecular structure (Figure 33) upon

annealing. The precursor flow rate did not appear to have a significant effect on film

properties on its own, but rather its effect was coupled with the effect of the plasma

power, since it was the concentration of ionized reactive species in the reactor that

affected the film structure.

90

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As-depositPost-anneal

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)

As-depositPost-anneal

Figure 34: (a) Siloxane and (b) organic ranges of FTIR spectra of as-

deposit and post-anneal VDEMS film deposited using 5 sccm VDEMS flowrate, 20 W plasma power, and 25 °C substrate temperature. A major change in the silicon bonding environments and a significant loss of organic content can be seen in the films after annealing.

High substrate temperature (200 °C) yielded films that had fairly low refractive

indices as-deposit and were well cross-linked, which was shown by the fact that neither

the thicknesses nor refractive indices changed upon annealing (Table 6). The cross-

linking in the films appeared to have taken place at the expense of organic moieties,

which agreed with the observations made by Gates et al. [42]. Substrate temperature

appeared to have a greater effect on film properties than either plasma power or precursor

flowrate, as both films deposited at 200 °C exhibited no change in molecular structure

(Figures 31 and 32), refractive index values, or film thickness. Depositing films at high

substrate temperatures was essentially similar to annealing them, since in both cases the

films have a molecular structure made-up of mainly long siloxane chains (Figures 30 and

31), which is dense, due to the loss of more volatile organic groups. This would explain

why no changes were seen in the films deposited at 200 °C after annealing.

91

Judging by the refractive indices, the VDEMS film deposited at 100 W, 1 sccm,

and 25 °C appeared to be the optimal film due to its low post-anneal refractive index and

good thermal stability. It would be possible to tolerate some, yet minimal, thickness loss,

if it is coupled with significantly better electrical and optical film properties.

4.1.2 Trivinyltrimethylcyclotrisiloxane Depositions

During preliminary depositions V3D3 proved to be a difficult precursor to plasma

polymerize as it appeared to have a narrow deposition space. Flowrates, reactor

pressures or plasma powers that were too high caused gas phase reactions to take place

and powder to form. This could be explained by the reactivity of the vinyl side chains,

which do not require much energy to polymerize, approximately 256 kJ/mol [46]. High

flowrates and reactor pressures yield a high precursor concentration in the reactor, and, in

turn, since the vinyl side-chains are so easily activated, a high concentration of ionized

species, which react at random in the gas phase before adsorbing onto the substrate.

Also, high plasma power causes a high concentration of ionized species in the reactor

compared to the concentration of the precursor molecules. Additionally, deposition

conditions that were too aggressive for the delicate functionality of the precursor, e.g.

high precursor flowrate, reactor pressure, plasma power, and high substrate temperature,

resulted in films with only minimal retention of cyclic structure. This could be explained

by the strained configuration of the silicon and oxygen atoms in the V3D3 ring structure:

even at low power depositions loss of precursor structure would be possible, as shown by

the wide siloxane band in Figure 35, which suggests heterogeneity of bonding

environments. Thus, mildest possible plasma conditions were used for the deposition of

92

V3D3 films: 1sccm V3D3 flowrate, 10 sccm argon flowrate, reactor pressure of 500

mTorr, substrate temperature of 25 °C, and 10 W plasma power. A small amount, 0.5

sccm, of t-butyl peroxide (tBPO) initiator was added to the inlet flow into the reactor to

activate the vinyl side chains on V3D3 precursor molecules and to help start the

polymerization reactions. These conditions allowed cyclic structure retention

demonstrated by the Si-O-Si band in Figure 35, the low wavenumber peak of which was

centered around 1010 cm-1, which is typically assigned to cyclotrisiloxane structures [41,

42]. The peak in the 1280-1250 cm-1 range pointed to incorporation of cyclic structure

into the films: it was centered at around 1270-1260 cm-1 indicating di- and tri-substituted

silicon atoms. While the siloxane band had an intense peak occurring at 1015 cm-1 it also

had a shoulder of significant intensity at higher wavenumbers, which suggested

heterogeneity of siloxane bonding environments.

93

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Figure 35: Segment of FTIR spectrum of PECVD V3D3 film

deposited at mild plasma conditions. The siloxane peak at 1015 cm-1 indicates retention of cyclic structure. The shoulder at higher wavenumbers indicates heterogeneity of siloxane bonding environments.

The siloxane band between 1200 and 1000 cm-1, illustrated in Figure 35, was

deconvoluted into five separate peaks to quantify the contributions of different types of

Si-O-Si bonding environments. The deconvolution, in Figure 36, shows that the major

contribution to the siloxane band was from cyclic Si-O-Si groups in the film; the peak at

1016 cm-1 accounted for 40 % of the total area under the curve. The second greatest

contribution was from long-chain siloxane groups, as indicated by peaks at 1050 and

1080 cm-1, which accounted for 38 % of the total area under the Si-O-Si band. Short-

chain siloxanes, showing up in the spectrum as peaks centered at 1129 and 1179 cm-1,

added only a small contribution to the intensity of the siloxane peak: 22 % of the total

94

area under the peak. So while the mild plasma conditions allowed for a sizable retention

of cyclic siloxane structures in the film, still more siloxane groups were present as linear

chains, both long and short, suggesting that even milder deposition conditions would

have to be utilized, to obtain a film with intrinsic porosity made up predominantly of

siloxane ring structures. It was most likely this breakdown of the cyclic structure that

lead to a fairly high refractive index of 1.482 ± 0.001 of the film described above.

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PECVD V3D3FitCyclicLong-Chain Long-Chain w/ OH or CH3Short-ChainShort-Chain w/ OH or CH3

1015.62 cm-1

40.3 %

1049.55 cm-1

15.9 %

1079.86 cm-1

22.0 %

1129.37 cm-1

18.36 %

1179.07 cm-1

3.3 %

Total Area = 3.627

Figure 36: Deconvolution of the siloxane peak of the PECVD V3D3 film spectrum. The greatest contribution to the peak area is from the cyclic siloxane peak at 1015 cm-1. However, more than half of the area under the peak is from various linear siloxane chains.

The V3D3 film was annealed for 90 minutes at 400 °C under nitrogen

atmosphere. The refractive index of the film decreased significantly to a new value of

1.422 ± 0.001. The decrease was likely due to the decline in the polarizability of the film

95

as a result of reacting of dangling bonds, as well as cross-linking that takes place between

hydroxyl groups with the elimination of water molecules that took place during the high

temperature annealing step. The decrease in the hydroxyl composition is demonstrated

by the loss of the peak between 3600 and 3100 cm-1 in the post-anneal film spectrum in

Figure 37.

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As deposit

Post-anneal

Figure 37: FTIR spectra of the PECVD V3D3 film, deposited at mild plasma conditions, before and after annealing. The main difference in the post-anneal spectrum is the disappearance of the hydroxyl peak, 3600-3100 cm-1, likely due to cross-linking that occurred between proximal hydroxyl groups during the annealing step with elimination of water molecules. Additionally, some cross-linking took place with elimination of methyl groups, as evident from the organic peak, 3100-2800 cm-1, which decreased in intensity in the post-anneal spectrum.

96

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As deposit

Post-anneal

Figure 38: Segment of FTIR spectra of the PECVD V3D3 film,

deposited at mild plasma conditions, before and after annealing. The shift in the siloxane peak towards higher wavenumbers in the post-anneal film spectrum points to the loss of some cyclic siloxane structure in the film. The shift in the peak between 1270 and 1250 cm-1 towards higher wavenumbers indicates cross-linking occurring in the film at high temperatures.

However, the drop off in the refractive index value was accompanied by 20 %

thickness loss upon annealing. The post-anneal film spectrum varied only slightly from

that of the as-deposit film, as can be seen in Figure 38. The peak intensity was nearly the

same in the as-deposit and post-anneal film spectra, which would suggest that no material

was lost with annealing and, thus, would conflict with the observed 20 % thickness loss

in the film. The most significant difference between the two spectra is the position of the

siloxane peak of greatest intensity, ~1015 cm-1, which can be attributed to cyclic siloxane

groups. In the post-anneal film spectrum it is shifted towards higher wavenumbers,

97

indicating loss of intrinsic pores. The oxygen substitution peak in the 1280-1250 cm-1

region also changed upon annealing: it shifted towards higher wavenumbers with only a

small shoulder at ~1260 cm-1. This change was consistent with the loss of the cyclic

structure in the film and reflected further cross-linking having taken place upon

annealing. The changes in the spectrum pointing to loss of cyclic siloxanes explain the

observed thickness loss in the film upon annealing: while no significant amount of

material was lost, the film densified as a result of the collapse of the intrinsic pores and

formation of additional cross-links.

The deconvolution of the siloxane band of the annealed film spectrum in Figure

38 shows the contributions of different types of Si-O-Si bonding environments to the

total Si-O-Si peak. The total area under the band is 3.485, only slightly lower than the

3.627 it was for the as-deposit spectrum. It can be seen that the peak indicating cyclic Si-

O-Si groups at ~1012 cm-1 is no longer a major contribution: it accounts for 19 % of the

total area under the peak, as opposed to 40 % prior to annealing. Linear long-chain

siloxane peaks account for 54 %, and the short-chain peaks contribute 26 % of total area.

The negligible decrease in total area along with the significant decrease in the percentage

of area contributed by the cyclic siloxane peak points to the fact that the thickness lost

upon annealing was due to the loss of Si-O-Si ring structure in the film, which caused the

intrinsic pores in the film to collapse.

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PECVD V3D3FitCyclicLong-Chain Long-Chain w/ OH or CH3Short-ChainShort-Chain w/ OH or CH3

1011.65 cm-1

19.4 %

1072.40 cm-1

30.3 %

1124.18 cm-1

24.2 %

1164.77 cm-1

2.2 %

Total Area = 3.485 1034.53 cm-1

23.9 %

Figure 39: Deconvolution of the siloxane peak of the post-anneal

PECVD V3D3 film spectrum. The total area under the peak remained nearly unchanged post-anneal. The contribution of the cyclic siloxane peak to the total area under the peak decreased by more than half upon anneal indicating a loss of intrinsic porosity during the high-temperature step.

In order to ascertain if deposition conditions favorable for obtaining VDEMS

films with low refractive indices would also yield a V3D3 with desired properties a

V3D3 film was deposited at a high substrate temperature of 200 °C. The V3D3 flowrate

was 2.5 sccm, tBPO flowrate was kept to a minimal 0.5 sccm, argon flowrate was 10

sccm, reactor pressure of 350 mTorr, and 20 W plasma power was used. The film was

later annealed for 90 minutes at 400 °C. Both as-deposit and post-anneal film spectra are

presented in Figure 40.

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Post-anneal

Figure 40: Segment of FTIR spectra of the PECVD V3D3 film,

deposited at high substrate temperature, before and after annealing. The siloxane peak occurs predominantly at 1035 cm-1 pointing to long-chain Si-O-Si structures within the film. The peaks for as-deposit and post-anneal films do not differ significantly as the film structure does not change greatly upon annealing due to the formation of well cross-linked films at high substrate temperatures. The oxygen substitution peak centered at 1270 cm-1 with a shoulder at 1260 cm-1 also indicates a cross-linked, as opposed to linear, Si-O-Si structure in the film.

It can be seen in the as-deposit film spectrum that the siloxane band has a peak of

highest intensity at ~1035 cm-1, indicating linear long-chain Si-O-Si structure in the film.

The oxygen substitution peak occurring at 1270 cm-1 with only a small shoulder at 1260

cm-1 also points to the loss of V3D3 intrinsic structure, in which each silicon atom has

bonds to only two oxygen atoms, and to the formation of cross-linked structure with tri-

substituted silicon atoms. The deconvolution of the silicon band in Figure 41 shows that

100

only about 10 % of the area under the peak is due to the contribution from the cyclic

siloxane peak at around 1020 cm-1. This loss of intrinsic structure is due to abundance of

energy in the reactor not only from plasma, but also thermal energy from having a heated

substrate, which breaks up the strained V3D3 ring structures easily. On the other hand,

the high substrate temperature causes very slow rates of adsorption onto the substrate,

and hence, slow deposition rates, thus resulting in a more thorough, ordered deposition,

without an abundance of dangling bonds. This is the reason for the fairly low refractive

index of the film: 1.425 ± 0.005.

The comparison between the as-deposit and post-anneal spectra in Figure 40 as

well as the post-anneal siloxane peak deconvolution in Figure 42 suggest that further loss

of cyclic structure took place upon annealing, as the intensity of the peak at ~1020 cm-1

decreased. The loss of cyclic structure appeared to be coupled with formation of Si-O-Si

linear short-chains, as the peaks at around 1130 and 1175 cm-1 increased in intensity, as

illustrated in Figure 40. While the structure of the film changed upon annealing there

was no material lost from the film, which is apparent from the area under the siloxane

peak, which is the same for the film before and after the deposition, shown in Figures 41

and 42, respectively. The thickness retention of the film after the high temperature

annealing step was ~98 %. The refractive index of the annealed film was 1.420 ± 0.002,

which, though is a low value, can still be obtained using non-cyclic precursors such as

VDEMS. The intrinsic porosity of the V3D3 precursor molecules was not taken

advantage of at these deposition conditions. Deposition conditions would need to be

identified, which would permit depositions in which the cyclic molecules both retained

their shape and were tightly cross-linked into the film matrix, so as to prevent the

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siloxane rings from opening up. This could possibly be achieved through depositions of

V3D3 along with a linear siloxane precursor, which would act as a spacer and alleviate

strain put on the film by trying to pack six-membered already-strained V3D3 rings

together.

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1020.26 cm-1

10.1 %

1035.04 cm-1

28.4 %

1074.71 cm-1

34.2 %

1133.32 cm-1

20.3 %

1176.39 cm-1

7.0 %

Total Area = 6.822

Figure 41: Deconvolution of the siloxane peak of the as-deposit

PECVD V3D3 film deposited at high substrate temperature. The biggest contribution to the area under the curve, ~ 63 %, is from long-chain siloxane peaks. The contribution from the cyclic peak is minimal, ~10 %, indicating the loss of cyclic precursor structure during deposition due to abundance of thermal energy from substrate heated to 200 °C.

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1017.80 cm-1

3.7 %

1033.32 cm-1

31.7 %

10741.32 cm-1

30.1 %

1123.05 cm-1

26.0 %

1172.15 cm-1

8.5 %

Total Area = 6.819

Figure 42: Deconvolution of the siloxane peak of the post-anneal

PECVD V3D3 film deposited at high substrate temperature. The total area under the peak remained unchanged post-anneal due to the well cross-linked structure of the film. The contribution of the cyclic siloxane peak to the total area under the peak decreased by 6.4 % upon annealing indicating a loss of intrinsic porosity. The contribution of the short-chain peaks increased by 7.2 %, indicating that the Si-O-Si rings open upon annealing to form short-chain structures.

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Figure 43: Comparison of FTIR spectra of V3D3-only films deposited at low (25 °C) and high (200 °C) substrate temperatures. The low substrate temperature film exhibited a greater retention of cyclic groups by having the greatest intensity of the siloxane peak at 1015 cm-1. The high temperature film was comprised primarily of long siloxane chains, since the peak occurred at around 1035 cm-1.

Comparing the spectra of the PECVD V3D3-only films deposited at low and high

substrate temperatures, presented in Figure 43, it could be seen that there was a trade-off

between cyclic structure retention and the thermal stability of the film. The film

deposited at low substrate temperature retained more of the cyclic siloxane structure

judging by the position of the siloxane peak of most intensity, which was centered at

lower wavenumbers (~1020 cm-1). However, the same mild conditions that aided in

retention of precursor structure, resulted in the deposition of a poorly cross-linked film,

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which underwent 20 % thickness loss upon annealing. The film deposited at high

substrate temperature, on the other hand, was mainly comprised of linear siloxane chains,

judging from the siloxane peak position at ~1035 cm-1, which was due to precursor

fragmentation brought on by the abundance of plasma as well as thermal energy in the

reactor. However, deposition at high substrate temperatures yielded a film structure that

is essentially similar to the structure of post-anneal low substrate temperature films: small

amount of cyclic groups, mainly made up of long-chain siloxanes, with some short-chain

groups. This could explain why the refractive indices of the post-anneal low temperature

film and the as-deposit high temperature film are so similar: 1.422 ± 0.001 and 1.425 ±

0.005, respectively.

4.1.3 Trivinyltrimethylcyclotrisiloxane-Vinylmethyldiethoxysilane

Depositions

The V3D3 film deposited at low substrate temperature (25 °C) had a low post-

anneal refractive index of 1.422 ± 0.001 but also underwent a ~20 % thickness loss upon

annealing. The film deposited at high substrate temperature (200 °C) was more

successful: a low post-anneal refractive index of 1.420 ± 0.002 was not accompanied by

any noticeable thickness loss as measured by spectroscopic ellipsometry. However,

comparable or even lower refractive indices were obtained for VDEMS-only films, thus

suggesting that full advantage is not being taken of the V3D3 intrinsic porosity. This

could be due to the fact that the six-membered siloxane ring structure is inherently

strained and becomes even more so when several molecules are cross-linked together.

105

What could alleviate this problem would be a linear spacer, which would lessen the strain

as V3D3 molecules would not have to be packed so closely together.

Since both V3D3 and VDEMS films deposited at high substrate temperature and

low plasma power had fairly low refractive indices and minimal thickness loss, a

combination V3D3-VDEMS film was deposited at a substrate temperature of 200 °C and

20 W power. The V3D3 flowrate was 2.5 sccm, the VDEMS flowrate 5 sccm, tBPO was

used at 0.5 sccm flowrate, and Ar flowrate was 10 sccm. Reactor pressure was 350

mTorr during the deposition.

In the comparison of the as-deposit V3D3 and V3D3-VDEMS film spectra in

Figure 44 it can be seen that the siloxane peak between 1200 and 950 cm-1 appears to be

shifted slightly towards lower wavenumbers in the V3D3-VDEMS film spectrum. This

is likely due to greater contribution of the cyclic peak, suggesting that using spacer

molecules alleviates the strain on the V3D3 molecules in the film, and allows for a

greater degree of ring structure retention in the films. The oxygen substitution peak,

between 1270 and 1250 cm-1, in the V3D3-VDEMS film spectrum occurs at slightly

lower wavenumbers and has a shoulder of greater intensity at 1260 cm-1 indicating either

the incorporation of VDEMS molecules and also possibly a greater extent of cyclic

siloxane structure in the film, in both of which silicon atoms have bonds to two oxygen

atoms. The refractive index of the as-deposit V3D3-VDEMS film was the same as that

of the V3D3-only film, within error: 1.432 ± 0.002, as opposed to 1.425 ± 0.005. But it

was still lower than 1.435 ± 0.002, the refractive index of the VDEMS-only film obtained

at similar deposition conditions. This can be due to the fact that the V3D3-only film has

a greater amount of long siloxane chains, which result from the opening of V3D3 rings in

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plasma. So, the any possible additional retention of cyclic precursor structure is offset by

the incorporation of small VDEMS molecules that increase the film density.

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Figure 44: Comparison of PECVD V3D3-only and V3D3-VDEMS film spectra, deposited at the same conditions. The siloxane peak in the V3D3-VDEMS film spectrum occurs at a lower wavenumber signifying better cyclic structure retention in the film with a spacer. The oxygen substitution peak also shows up at slightly lower wavenumbers and has a bigger shoulder at 1260 cm-1 in the V3D3-VDEMS film spectrum.

From the peak deconvolution in Figure 41 it can be seen that ~63 % of the area

under the siloxane peak in the V3D3-only film spectrum is due to the contribution of

long-chain Si-O-Si groups. Figure 45, however, shows that the long-chain siloxane peaks

make up only about 57 % of the area under the overall siloxane peak in the V3D3-

VDEMS film spectrum. The long linear siloxane chains could act as linear spacers

between the siloxane rings preserved in the film, thus yielding a film with both intrinsic

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and induced porosity. Thus, it can be concluded that the VDEMS molecules, while aid in

the preservation of the V3D3 cyclic structure, do not contribute to the porosity of the film

themselves.

The deconvolution of the siloxane peak from the post-anneal spectrum presented

in Figure 46 shows that the cyclic peak at ~1015 cm-1 decreases by about a quarter upon

annealing: from 15.3 % to 11.8 % of the total area under the peak. This decrease is

significantly less than that evident in the spectrum of the V3D3-only film upon annealing,

which decreased by about two thirds from 10.1 % to 3.7 % of the total area under the

peak. This points to better retention of cyclic structure in films with spacer molecules, as

they decrease the strain on V3D3 molecules, and thus, allow for a better cross-linked

more stable films. Upon annealing the refractive index of the film did not change greatly:

the post-anneal value was 1.422 ± 0.005. This was marginally higher than the 1.420 ±

0.002 refractive index value of the post-anneal V3D3-only film, but still quite a bit lower

than the 1.435 ± 0.001 value of the post-anneal VDEMS-only film. This suggests that the

addition of VDEMS molecules as spacers did not benefit the porosity of the film: while

more of the V3D3 cyclic structure was retained in the final film, the small VDEMS

molecules caused the film density to increase. Utilizing linear spacer molecules with

longer chain lengths could help achieve films with lower refractive indices. Linear

spacers would have two-fold contribution to the porosity of the films: they would

contribute to the induced porosity of the film as well as space V3D3 molecules apart so

as to alleviate the strain, and thus, ensure intrinsic porosity by helping the siloxane rings

stay intact.

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1015.42 cm-1

15.3 %

1034.52 cm-1

26.4 % 1074.91 cm-1

30.5 %

1131.00 cm-1

23.5 %

1174.27 cm-1

4.3 %

Total Area = 5.983

Figure 45: Deconvolution of the siloxane peak of the as-deposit

PECVD V3D3–VDEMS film deposited at high substrate temperature. There is a significant contribution to the area under the curve, ~ 15 %, is from the cyclic siloxane peak, indicating greater retention of cyclic structure than in the V3D3-only film, likely due to VDEMS molecules alleviating the strain on the V3D3 molecules.

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1017.47 cm-1

11.8 %

1036.34 cm-1

33.0 %

1087.04 cm-1

26.8 %1131.50 cm-1

27.5 %

1177.70 cm-1

0.9 %

Total Area = 6.183

Figure 46: Deconvolution of the siloxane peak of the post-anneal

PECVD V3D3-VDEMS film deposited at high substrate temperature. The total area under the peak remained unchanged due to the well cross-linked film structure. The lesser decrease in the percentage of the cyclic peak intensity compared to V3D3-only film was due to a more stable structure as result of VDEMS spacer molecules.

4.1.4 Trivinyltrimethylcyclotrisiloxane-Spacer Depositions

To test the hypothesis that linear spacer molecules would contribute to induced

porosity as well as aid in preservation of intrinsic porosity V3D3-spacer films were

deposited using linear siloxane spacer molecules. Spacer precursors with end-capping

vinyl bonds and varying siloxane chain lengths were utilized: divinyldimethylsilane (S1),

divinyltetramethyldisiloxane (S3), and divinylhexamethyltrisiloxane (S5). According to

the hypothesis the longer the siloxane chain of the spacer molecule, the lower resulting

film density, and the lower the refractive index and the dielectric constant of the film. All

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films were deposited using 1 sccm V3D3 flowrate, 1.5 sccm spacer flowrate, minimal

flowrate of tBPO of 0.5 sccm, Ar flowrate of 50 sccm, and a reactor pressure of 500

mTorr. The stage temperatures of 25 °C and 200 °C, and deposition powers of 10 and 20

W were used. V3D3-S3 and V3D3-S5 films were annealed in nitrogen atmosphere at

400 °C for 90 minutes.

Table 7: Optical properties and thermal stability data for the V3D3-spacer films deposited at varying substrate temperatures and deposition powers

Conditions Spacer Refractive Index (AD)

Refractive Index (PA)

Thickness Loss (%)

25 °C and 20 W S1 1.495 ± 0.001 S3 1.482 ± 0.001 1.462 ± 0.001 22.2 ± 1.5 S5 1.476 ± 0.003 1.457 ± 0.001 16.6 ± 0.3

25 °C and 10 W S1 1.493 ± 0.002 S3 1.482 ± 0.001 1.464 ± 0.001 30.2 ± 0.5 S5 1.476 ± 0.003 1.460 ± 0.006 31.9 ± 1.3

200 °C and 20 W S1 1.494 ± 0.001 S3 1.448 ± 0.001 1.440 ± 0.001 0.2 ± 2.2 S5 1.446 ± 0.001 1.445 ± 0.002 0.6 ± 1.3

200 °C and 10 W S1 1.476 ± 0.003 S3 1.433 ± 0.001 1.422 ± 0.001 0.8 ± 0.5 S5 1.433 ± 0.001 1.427 ± 0.001 2.1 ± 3.3

It can be seen in Table 7 that for each set of deposition conditions as-deposit V3D3-S5

films had the lowest refractive index values of all V3D3-spacer films, thus confirming

that the spacer chain length has an effect on the molecular structure of the film.

However, the chain length of the spacer molecules had a smaller effect on the optical

properties of the as-deposit films than prognosticated: the difference in refractive indices

of V3D3-S3 and V3D3-S5 films was statistically insignificant in two out of four

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conditions. The refractive index values of the low substrate temperature films were

higher than those of the high substrate temperature films as-deposit, which was likely due

to their poorly cross-linked structure, with abundance of dangling bonds, which increased

the polarizability of the films. The low substrate temperature films also exhibited a much

more significant decrease in refractive index values post-anneal, which was coupled with

thickness losses of as high as 30 %. The post-anneal reduction in the refractive index

value was due to reactions among dangling bonds, further cross-linking and formation of

more ordered film structure that took place during the high temperature annealing step.

The lowest refractive indices for each type of film were obtained when the films were

deposited at high substrate temperature and low power. Deposition rates at these

conditions were slow, which allowed for more thorough cross-linking, minimizing the

number of dangling bonds in the film, and resulted in films with highly ordered structure.

Figure 47 shows a comparison of V3D3-S3 films deposited at 25 and 200 °C. The ratio

of cyclic and long-chain peaks to short-chain peaks in the low substrate temperature film

spectrum is substantially lower than in the spectrum of the high substrate temperature

film. This explains the significant difference between the refractive indices of the two

films, since it is siloxane ring structure and long chains that reduce the density of the

film, and thus, its refractive index. The poor cross-linking in the low substrate

temperature film is manifested by the oxygen substitution peak between 1270 and 1250

cm-1, and which in this case appears between 1260 and 1250 cm-1 with only a slight

shoulder at 1270 cm-1, indicating poorly cross-linked structure mainly comprised of

mono- and di-substituted silicon atoms. Conversely, the V3D3-S3 film deposited at 200

°C has an oxygen substitution peak that shows up as a doublet at 1270 and 1260 cm-1

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with a small shoulder at 1250 cm-1, pointing to its better cross-linked more ordered

structure.

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200 deg C, 10 W

Figure 47: Segment of FTIR spectra of the PECVD V3D3-S3 films, deposited at 25 °C, 10 W and 200°C, 10 W.

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V3D3V3D3+S1V3D3+S3V3D3+S5

Figure 48: Comparison of V3D3-only and V3D3-spacer film spectra deposited at 200 °C and 10 W. The spectra suggest that the V3D3+S3 film has the greatest cyclic content of the three films with spacers, which would explain why it has the lowest refractive index value of the three. The V3D3-only film has a refractive index value lower than the three V3D3-spacer films, which can be explained by it having a smaller ratio between cyclic and long-chain to short-chain peak intensities.

The as-deposit spectra of the three V3D3-spacer films deposited at 200 °C and 10

W are presented in Figure 48 and compared with the spectrum of the V3D3-only film

deposited at high stage temperature. It can be seen that the siloxane peak in all three

V3D3-spacer film spectra is at the greatest intensity at low wavenumbers, ~1020 cm-1,

indicating considerable contribution from a cyclic peak. Also, the V3D3-S5 film, the

spacer in which had the longest siloxane chain, appeared to have the greatest cyclic and

long-chain peak to short-chain peak ratio, which must have been the reason for its low

114

refractive index of 1.433 ± 0.001. However, the V3D3-S3 film that had a shorter spacer

length also had an as-deposit refractive index value of 1.433 ± 0.001. This could be

explained by a greater number of V3D3 cyclic structures retained in this film, which is

evident from the higher intensity of the siloxane peak at lower wavenumbers. While

fairly low refractive indices were achieved by depositing V3D3-spacer films at high

substrate temperature, the refractive index of the as-deposit V3D3-only film at high

substrate temperature was still lower at 1.425 ± 0.005. This could be explained by the

much greater ratio of the cyclic and long-chain peaks to the short-chain siloxane peaks in

the V3D3-only films spectrum than in any of the V3D3-spacer film spectra. This

suggests that some densification takes place in the V3D3-spacer films as a result of

monomer fragmentation. In case of the V3D3-S1 films, the addition of the spacer did not

improve the optical performance of the film, but in fact impaired it, as the small size of

the spacer caused the film density to increase, similar to the V3D3-VDEMS film.

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1013.18 cm-1

15.1 %

1030.62 cm-1

26.7 %

1079.57 cm-1

34.4 %

1137.04 cm-1

20.1 %

1176.88 cm-1

3.7 %

Total Area = 5.841

Figure 49: Deconvolution of the siloxane peak of the as-deposit PECVD V3D3–S3 film deposited at 200 °C and 10 W. There was a significant contribution to the area under the curve, ~ 15 %, by the cyclic siloxane peak, indicating significant retention of cyclic structure.

While the V3D3-S3 and V3D3-S5 films deposited at 200 °C and 10 W had the

same low refractive index as-deposit, 1.433 ± 0.001, the V3D3-S3 film had a slightly

lower refractive index post-anneal than the V3D3-S5 film: 1.422 ± 0.001, as opposed to

1.427 ± 0.001. This could be explained by the fact that the V3D3-S3 film had a greater

siloxane ring structure and long-chain content as-deposit than the V3D3-S5, which can be

seen in the comparison of the two spectra in Figure 48. The deconvolution of the

siloxane peak of the as-deposit V3D3-S3 film in Figure 49 showed that the contribution

from the cyclic peak was ~15 % and the contribution from two long-chain siloxane peaks

added up to ~61 %. Comparing this deconvolution to that of the post-anneal V3D3-S3

116

film in Figure 50 it could be seen that the area under the peak stayed nearly the same, as

did the contributions from the different siloxane peaks: the cyclic peak accounted for ~15

% of the area under the peak, while the long-chain peaks contributed ~57 %.

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1019.85 cm-1

15.6 %

1034.51 cm-1

33.0 %

1078.76 cm-1

23.9 %

1113.41 cm-1

20.1 %

1159.94 cm-1

7.6 %

Total Area = 5.271

Figure 50: Deconvolution of the siloxane peak of the post-anneal PECVD V3D3-S3 film deposited at 200 °C and 10 W. The total area under the peak decreased slightly post-anneal, albeit without any significant loss of cyclic structure.

Thus, the lowest refractive index obtained for a V3D3-spacer film was 1.422 ±

0.001: that of the annealed V3D3-S3 film deposited at 200 °C and 10 W. The same

refractive index value was obtained for the post-anneal V3D3-VDEMS film deposited at

200 °C and 20 W. The lowest refractive index was that of the annealed V3D3-only film,

which was deposited at 200 °C and 20 W: 1.420 ± 0.002. This suggests that the addition

of spacer co-monomers did not benefit the porosity of the film: even at low plasma

117

powers there was an abundance of thermal energy in the reactor from the high substrate

temperature to cause monomer fragmentation of both the V3D3 cyclic structures and the

spacer siloxane chains to increase film density. Also, it appeared that the addition of

small spacer molecules, such as S1, cause the film density to increase. Thus, a low

energy CVD technique needed to be utilized to ensure precursor structure retention in the

final film.

4.2 Initiated CVD

The goal of depositing organosilicon thin films with intrinsic porosity using

V3D3 as a precursor was only partly achieved when PECVD was used as a deposition

technique. This was hypothesized to be due to the fact that the plasma, even at low

powers, provided such an abundance of energy to the monomer species in the reactor that

the delicate precursor functionality was not retained in its entirety in the deposited film.

Additionally, the continuous plasma subjected the growing film to constant ion

bombardment resulting in even further loss of strained V3D3 ring configuration. This

was believed to be the reason why the lowest refractive index obtained for an as-deposit

PECVD V3D3 film was 1.425 ± 0.005, and 1.433 ± 0.001 for an as-deposit V3D3-spacer

film, while a refractive index value that was just as low, 1.430, was obtained for a

VDEMS film, where the precursor had no intrinsic porosity. Thus, a better, milder

deposition method was necessary for V3D3 films to retain as much of their intrinsic pores

as possible in order to yield films with lower refractive indices. iCVD was chosen as a

gentler alternative to PECVD to avoid plasma excitation and damage to the growing film

as it provides only enough thermal energy to cleave the easily activated thermally labile

118

bond in the initiator species, while keeping the monomer structure intact. Additionally,

since nominal energy is applied, iCVD provides more control over the polymerization

pathways that can take place.

For all iCVD depositions 80/20 Nichrome filament was used positioned in a

parallel array, so as to guarantee uniform temperature. The filament array position was

fixed at 2.5 cm above the substrate. The initial depositions focused on V3D3 film

structure and deposition rate as a function of filament temperature, thus, depositions were

conducted with filament temperatures ranging from 150 to 350 °C. tert-Butyl peroxide

was used as thermally labile initiator species. Its flowrate was kept low at 1 sccm, so as

to ensure the dominance of propagation reactions during polymerization. V3D3 flowrate

was kept at 5 sccm, the reactor pressure was 250 mTorr, and the substrate was kept at 25

°C by back-side cooling. At filament temperatures of less than ~200 °C no deposition

was apparent. Also, no deposition occurred at any filament temperature without the

presence of tBPO. At filament temperatures of 200, 250 and 350 °C the deposition rates

were 22.27 ± 3.70, 32.83 ± 2.04, and 71.78 ± 8.48 nm/min respectively, suggesting a

dependence of deposition rate on filament temperature, which was consistent with

literature [26]. At higher temperatures the tBPO species were more likely to form

initiating radicals, and thus initiate more polymer chains, resulting in faster deposition

rates. The effect of filament temperature on film structure could be construed from the

siloxane peak segment of FTIR spectra of the V3D3 films deposited at 200 and 350 °C,

presented in Figure 51. The siloxane peak of the film deposited at 200 °C filament

temperature was narrow and singular, occurring at around 1010 cm-1, pointing to almost

complete retention of the cyclic structure of the precursor. On the other hand, the

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siloxane peak of the film deposited at 350 °C filament temperature was wider, shifted to a

slightly higher wavenumber, pointing to heterogeneity of bonding environments, with a

shoulder at around 1085 cm-1, which identified with long-chain linear siloxanes with

terminal methyl or hydroxyl groups.

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Filament T ~ 200 °C

Figure 51: Segment of FTIR spectra of iCVD V3D3-only films deposited at filament temperatures of 200 and 350 °C. The siloxane peak of the film deposited at higher filament temperature shows a shift towards higher wavenumbers, suggesting some loss of cyclic structure, as well as a shoulder at higher wavenumbers, pointing to heterogeneity of bonding environments.

Judging from the deconvoluted spectra of both the films presented in Figures 52

and 54 it could be seen that in both cases the cyclic peak at around 1010 cm-1 was the

major contributor to the area under the peak. However, its contribution was greater in the

film deposited at lower filament temperature, of 200 °C, 83.4 %, as opposed to 70.2 % in

120

the spectra of the film deposited using a filament temperature of 350 °C. The spectrum

of the 350 °C filament temperature film also had significant contribution from siloxane

long-chains, 21.4 %, suggesting that chain opening must have taken place during the

depositions due to excess thermal energy.

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CyclicLong-Chain Long-Chain w/ OH or CH3Short-Chain

963.84 cm-1

5.0 %

1008.84 cm-1

83.4 %

1020.83 cm-1

1.0 %1068.5 cm-1

10.6 %

Total Area = 3.565

Figure 52: Siloxane peak deconvolution of the as-deposit iCVD V3D3-only film deposited at the filament temperature of 200 °C.

The refractive indices of the films also tracked with the filament temperatures:

1.454 ± 0.002, 1.453 ± 0.007, and 1.473 ± 0.005 for the films deposited at 200, 250 and

350 °C respectively. While the film deposited at a filament temperature of 200 °C did

not have the lowest refractive index, it was still low considering the range of error in the

refractive index measurement of the film deposited at 250 °C. Also, the two filament

temperatures, 200 and 250 °C, were within the error in filament temperature

121

measurement, which was ± 50 °C. Thus, refractive indices that were close in value

would be expected. What was notable was that the presence of linear siloxane groups in

the 350 °C V3D3 film, evident in the spectrum in Figure 53, resulted in a significant

difference in refractive indices. Hence, for all other iCVD film depositions the mildest

filament temperature of 200 °C was used to ensure minimal precursor fragmentation and

structure loss.

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965.43 cm-1

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70.2 %

1019.8 cm-1

1.2 %

1188.39 cm-1

21.4 %

1148.84 cm-1

5.7 %

Total Area = 4.774

Figure 53: Siloxane peak deconvolution of the as-deposit iCVD V3D3-only film deposited at the filament temperature of 350 °C.

Comparing the refractive indices of the as-deposit V3D3-only films deposited via

PECVD and iCVD techniques it had to be kept in mind that the PECVD film deposited at

200 °C substrate temperature had a significantly lower refractive index than the iCVD

film, 1.425 ± 0.005, as opposed to 1.454 ± 0.002, due to the fact that high substrate

122

temperature had the same effect on the growing film as annealing that was done after

deposition. Thus, for a more accurate comparison between the two techniques the

PECVD film that was deposited at a substrate temperature of 25 °C should have been the

one compared to the iCVD film. Thus, comparing the refractive index values of the

PECVD (low substrate temperature) and iCVD V3D3-only films, 1.482 ± 0.001 and

1.454 ± 0.002, respectively, it could be seen that the milder deposition conditions of the

iCVD technique yielded a film with a greater extent of intrinsic porosity because of better

precursor structure retention. The vast difference in the structure of the two films could

be seen in the comparison of the siloxane peaks of the FTIR spectra in Figure 54. The

areas under the two peaks were nearly the same: 3.627 for the PECVD film (from Figure

36) and 3.565 (from Figure 52). What was different, however, was the shape of the

peaks: the siloxane band in the iCVD film spectrum was narrower and more singular,

pointing to the film structure being made up predominantly of siloxane rings, while the

band in the spectrum of the PECVD film, though of less intensity, was wider, which

pointed to heterogeneity of bonding environments in the film. The comparison of the two

films indicated that iCVD was a more suitable technique for depositing films from

precursors with delicate functionalities. Since the iCVD films were not annealed it was

impossible to compare the optical properties of the post-anneal films across the different

CVD techniques.

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Figure 54: Comparison of FTIR spectrum segments of PECVD and iCVD V3D3-only films. Both films were deposited with minimal energy input: the PECVD film was deposited at a substrate temperature of 25 °C and 10 W power, iCVD film - at the filament temperature of 200 °C and a substrate temperature of 25°C.

Since it was observed from the FTIR spectrum that iCVD yields V3D3 thin films

with a great extent of intrinsic porosity, V3D3-spacer films were deposited in order to

extend film porosity by adding induced pores into the film. The same spacer precursors

were used in the iCVD films as in the PECVD films: S1, S3, and S5. All films were

deposited at a filament temperature of 200 °C and substrate temperature of 25 °C. In all

cases V3D3 flowrate was kept at 5 sccm, spacer flowrate was 7.5 sccm, tBPO flowrate of

1 sccm was used, and the reactor pressure was kept at 250 mTorr. The refractive index

values for each film are presented in Table 8.

124

Table 8: Optical properties for the V3D3-spacer films deposited using the iCVD technique

Spacer Refractive Index (AD)

none 1.454 ± 0.002

S1 1.481 ± 0.001

S3 1.472

S5 1.461 ± 0.002

It could be seen from Table 8 that for V3D3-spacer film depositions the V3D3-S5

films had the lowest refractive index values, thus showing again that the chain length of

the siloxane spacer had an effect on the optical properties of the film. However, it could

be seen from Table 8 that while refractive index values decreased for the V3D3-spacer

films as the spacer siloxane chain length increased, all of the V3D3-spacer films had

refractive indices that were higher than that of the V3D3-only film. This was the same

trend that was observed in PECVD V3D3-spacer films, the optical constants of which are

given in Table 9.

Table 9: Comparison of optical properties of the V3D3-spacer films deposited using the iCVD and PECVD techniques

Spacer Refractive Index (PECVD)

Refractive Index (iCVD)

None 1.482 ± 0.001 1.454 ± 0.002

S1 1.493 ± 0.002 1.481 ± 0.001

S3 1.482 ± 0.001 1.472

S5 1.476 ± 0.003 1.461 ± 0.002

125

While in the case of PECVD films this discrepancy in the optical properties could

be attributed to spacer chain fragmentation and film densification, these were not likely to

occur in the iCVD process, which was mild enough for structure retention of much more

delicate V3D3 molecules. This observation suggested that our hypothesis of linear

siloxane molecules with two end-capping vinyl groups arranging themselves in a fashion

of spacers among V3D3 molecules, as illustrated in Figure 27, was incorrect. It was

likely that in iCVD where the energy was not nearly as abundant as in PECVD and less

reaction pathways were available for the precursors. So while the spacers polymerized

through the vinyl bonds, the polymerization was likely similar to that of acrylates, so that

the linear part of the molecule acted as a pendant group. This would explain the observed

film densification and the increase in the refractive index with addition of the spacers to

V3D3 films.

Comparing the refractive index value of each film with that of the same V3D3-

spacer film deposited using the PECVD technique and the mildest conditions, seen in

Table 9, it could be seen that in fact iCVD films performed better and had significantly

lower refractive indices. Also, the comparison of iCVD and PECVD film spectra for

V3D3+S3 and V3D3+S5 films presented in Figures 55 and 56, showed that in both cases

the iCVD film spectra had narrower peaks suggesting better defined film structure with

only a few types of bonding environments in the film. The iCVD film spectra also

insinuated much better preservation of precursor structure. Both PECVD film spectra

had wider siloxane bands and significantly lower intensity at 1015 cm-1, where the cyclic

peak showed up. This yet again showed that iCVD was in fact a better method of

deposition of films using cyclic precursors with delicate structure, such as V3D3.

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iCVD V3D3+S3

Figure 55: Comparison of FTIR spectrum segments of PECVD and iCVD V3D3-S3 films deposited with minimal energy input.

127

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Figure 56: Comparison of FTIR spectrum segments of PECVD and iCVD V3D3-S5 films deposited with minimal energy input.

128

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V3D3-S1

V3D3-S3

V3D3-S5

Figure 57: iCVD V3D3-spacer films using S1, S3 and S5 deposited at the same conditions. The appearance of the peak at around 1120 cm-1 indicates the incorporation of linear siloxane chains into the film.

This densification could also be observed spectrally in Figure 57, where the

appearance of a peak at ~1120 cm-1 indicated the incorporation of linear siloxane short-

chains into the film, not the formation of long chains, which would result in the formation

of significant induced porosity in the film. Some references [41, 42] also identified the

peak at these wavenumbers as cage structure, which would fit better the hypothesis that

the spacer films polymerized similar to acrylates and the linear molecule acted as a

pendant group. It could also be seen in the spectra in Figure 57 that the longer the spacer

chain the less intense the siloxane peak was. This meant that there were also less cyclic

siloxane structures in these films, which would account for refractive indices that were

lower than expected.

129

Following the work of Lau and Gleason [52] the effect of PM/Psat ratio, the ratio of

the monomer gas pressure over its saturated vapor pressure, on V3D3 deposition rate and

film structure was investigated. It was shown by Lau and Gleason [52] that the main

difference between iCVD and PECVD wass where the reaction took place. Because of

the abundance of energy throughout the reactor in plasma depositions the reactants are

continuously activated as they move through the plasma, and as such, fragmentation,

recombination, and gas-phase polymerization are much more likely to occur.

Consequently, if the reactant flowrate is increased PECVD, an increase in deposition rate

can be seen, as oligomers grow in the gas phase, then diffuse to the surface and contribute

to the film. On the other hand, iCVD was shown to be mass-transfer limited, and the

precursor adsorption to the surface was shown to be the rate-limiting step, since the

activation is low-power and confined to a narrow filament. Virtually no gas-phase

polymerization was shown to occur, and in order for the film to grow, the species had to

make it to the surface and stick to it. Thus, the precursor flowrate made no difference,

only the precursor species adsorbed onto the surface add to the growth of the film.

PM/Psat thus was shown to be very important, as it determined the concentration of the

precursor species on the substrate surface.

All depositions were performed using a V3D3 flow rate of 5 sccm and a reactor

pressure of 250 mTorr, thus, keeping PM constant between all the runs. In order to vary

the PM/Psat ratio for each deposition the substrate temperature was varied in 5 °C

increments from 15 to 30 °C, thus, changing the Psat for each of the runs. All the rest of

the deposition variables were fixed: tBPO flowrate of 1 sccm was used, and the filament

wires were heated to 200 °C. As presented in Table 10, the deposition rate of the films

130

increased with increasing PM/Psat, pointing to monomer adsorption onto the substrate

being a primary factor in iCVD polymerization. The refractive indices of the four films

are statistically the same, with the exception of the film deposited at the highest PM/Psat

ratio of 0.77, which is higher and slightly outside the error margins of the other refractive

index values of the other films. Overall, the film deposited at an average PM/Psat value of

0.35 had the lowest refractive index with the smallest variations across the whole film,

and hence with the lowest error margin, and a reasonable deposition rate.

Table 10: Comparison of deposition rates and optical properties of iCVD V3D3-only films deposited using different PM/Psat ratios

Substrate Temperature (°C)

PM/Psat Deposition Rate (nm/min)

Refractive Index

15 0.77 32.12 ± 0.17 1.463 ± 0.004 20 0.52 27.18 ± 0.78 1.455 ± 0.004

25 0.35 20.21 ± 0.63 1.454 ± 0.002

30 0.23 14.98 ± 0.69 1.458 ± 0.007

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Pm/Psat=0.77, Stage T=15 ° C

Pm/Psat=0.52, Stage T=20 ° C

Pm/Psat=0.35, Stage T=25 ° C

Pm/Psat=0.23, Stage T=30 ° C

Figure 58: iCVD V3D3-only films using different Pm/Psat ratios controlled by changing the substrate temperature.

Spectrally the films also had only minor variations: the siloxane cyclic peak

shifted down by no more than a few wavenumbers as the PM/Psat value decreased. This

was likely due a more structured film formation at lower higher substrate temperatures.

Since it was shown by Lau and Gleason [52] and by our own depositions of

V3D3-only films that monomer adsorption onto the substrate surface and subsequently its

polymerization was dependent on its PM/Psat value it could be assumed that the

copolymerization of V3D3 and a spacer would be more successful and would yield better

results if the two monomers had PM/Psat ratios close in value. V3D3-S5 film was

deposited using a PM/Psat value of 0.42 for both of the precursors. The value was chosen

as it was easily obtained for both precursors at a substrate temperature of 15 °C, reactor

132

pressure of 250 mTorr and flowrates of 5 sccm for each monomer. tBPO flowrate of 1

sccm was used. The filament temperature was 200 °C. A V3D3-only film was deposited

at the same conditions for comparison. Argon gas flowrate of 5 sccm was used in order

to maintain the PM/Psat value as for the copolymer film. The deposition rate of the V3D3-

S5 film was more than double that of the V3D3-only film, 28.91 ± 0.44 nm/min and

12.42 ± 0.29 nm/min, respectively, indicating the deposition of the additional material in

case of the V3D3-S5 film. The comparison of the siloxane peak in the FTIR spectra of

the two films in Figure 59 showed the incorporation of linear siloxane groups into the

film with the addition of S5. The siloxane peak shifted only slightly towards

wavenumbers indicating no apparent loss of cyclic structure. The band broadened

indicating a higher intensity of long-chain peak at ~1035 cm-1 and CH3/OH-capped long-

chain peak at ~1085 cm-1. There was also some apparent increase in intensity around

1135 cm-1 indicating some incorporation of short-chain or cage-structure siloxane groups.

The incorporation of S5 linear chains into the film was also manifested by the increase of

the oxygen substitution peak, which increased in intensity at 1260 cm-1, which implied

the incorporation of di-substituted silicon atoms into the film structure.

133

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V3D3 Pm/Psat=0.42 T=15 °C

V3D3-S5 both Pm/Psat=0.42 T=15 °C

Figure 59: iCVD V3D3 and V3D3+S5 films deposited at the same conditions with the same Pm/Psat ratios.

The refractive index value exhibited a decrease with the addition of the spacer

species from 1.462 ± 0.002, for the V3D3-only film, to 1.453 ± 0.010, for the V3D3-S5

film, but the two values were statistically the same. However, this was an improvement

on the PECVD and previous iCVD films, where the addition of spacer caused the film to

densify and the refractive index to increase. So in this case of V3D3-S5 film where both

precursors had the same PM/Psat ratio during deposition the addition of spacer moieties

actually may have added some porosity to the films. It could be seen from Table 10 that

the PM/Psat ratio had some effect on the optical properties of the film: there may be a

window of optimal PM/Psat values that would allow better control over film composition.

134

The refractive index of the V3D3-S5 film deposited using the same PM/Psat ratio

for both precursors was also lower than that of the iCVD V3D3-S5 film deposited

without consideration for PM/Psat: 1.453 ± 0.010 and 1.461 ± 0.002 respectively.

However, it was difficult to judge the extent of improvement in optical properties due to

the large error margin in the value of the former film.

The comparison of FTIR spectra of PECVD (low substrate temperature and low

power) and iCVD (PM/Psat=0.42) V3D3-S5 films in Figure 60 pointed to a significant

difference between the two films. The iCVD film exhibited a siloxane peak of

significantly higher intensity as well as better retention of V3D3 structure and uniformity

and homogeneity of film makeup as evidenced by the position of the peak at around 1015

cm-1, and the fact that the peak was much narrower. The wider peak at higher

wavenumbers in the PECVD film spectrum pointed to loss of cyclic structure and

unsystematic bond formation owing to the excess of input energy during deposition. The

deconvolution (not shown) of both siloxane peaks from Figure 60 showed that the area

under the peaks was nearly the same: 3.951 for the iCVD film, versus 4.035 for the

PECVD. This suggested that the same amount of monomer was polymerized and that the

only difference between the two films was the bonding environment: the fact that, in once

case, the cyclic structure was retained and in the other it wasn’t, resulting in a denser film

with more heterogeneous bonding environments. The comparison of the refractive

indices of the films: 1.453 ± 0.010 for the iCVD film and 1.476 ± 0.003 for the PECVD

film, supported the conclusion that iCVD was a gentler technique that allowed

depositions of V3D3 films with intrinsic porosity, and, at right deposition conditions,

135

with the addition of linear co-monomers, depositions of V3D3-spacer films with added

induced porosity.

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Figure 60: Comparison of FTIR spectrum segments of PECVD (low substrate temperature and low power) and iCVD (PM/Psat=0.42) V3D3-S5 films. The area under the peaks was nearly the same, and the difference between the films was in the bonding environments present in each film.

Thus, iCVD was indeed shown to be a more suitable technique when the

deposition of V3D3 films was concerned, since it allowed for excellent retention of cyclic

structure and yielded films with considerably lower refractive indices than PECVD

technique at similar low energy conditions. The concentration of monomer on the

substrate surface, and, consequently, the PM/Psat ratio for each monomer proved to be of

importance for iCVD polymerizations, as suggested by Lau and Gleason [52, 53]. This

variable was decisive for co-polymerizations, where control over the amount of each

136

monomer in the final film was needed, due to the dependence of deposition rate of each

co-monomer on its respective PM/Psat. More work is necessary for optimization of

deposition conditions for V3D3-spacer co-polymers.

4.3

It has been shown in literature that surfaces can be biofunctionalized via protein

attachment to carbonyl groups on the surface. So one of the important characteristics for

our poly(acrylic acid) (pAA) films is the density of carboxylic acid COOH functional

groups on the film surface, since it will control the degree to which the film surface can

be biofunctionalized in the future. In order to obtain films with high carboxyl group

density the precursor structure needs to be retained in the final film and the

polymerization of the films needs to occur only along the vinyl bond of the precursor

acrylic acid (AA) molecules. Plasma enhanced CVD (PECVD) can often provide too

much excitation in the plasma, which results in the loss of precursor structure. Thus, low

plasma powers need to be utilized during deposition.

Deposition of Biocompatible Insulating Films

The effect of plasma power was investigated by depositing a series of pAA films

at constant deposition conditions: 1 sccm of AA, 1 sccm of tBPO, and reactor pressure of

350 mTorr. Four depositions were performed at plasma powers of 10, 20, 50 and 100 W.

It can be seen in the FTIR spectra of the films presented in Figure 61 that the intensity of

the carbonyl peak at 1760-1700 cm-1 decreased with increasing plasma power, indicating

a loss of carbonyl groups in the film. This was probably due to the fact that at greater

powers radical formation in plasma was enhanced, leading to greater degree of precursor

fractionation and deposition of more irregular films with high degree of cross-linking.

137

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100 W

50 W

20 W

10 W Increasing plasma power

Figure 61: Segment of FTIR spectra of pAA films deposited at

varying plasma powers. The decrease in the carbonyl peak intensity (1740-1700 cm-1) with increasing plasma power pointed to loss of precursor structure and cross-linking occurring in the film due to high concentration of radicals forming in the plasma.

This was confirmed by testing film stability in water. The films were soaked in

DI water for 1 hour. It can be seen in Figure 62 that the higher the plasma power, the

better the film thickness retention, which pointed to greater degree of cross-linking in the

film. Films deposited at plasma powers of 50 and 100 W showed nearly 100 % thickness

retention post-soak. While it is important for the films to be stable in aqueous

environments, this stability should not be enhanced at the expense of carboxyl groups in

the films, since these are the sites of protein tethering for biofunctionalization. Thus,

pAA films need to be deposited at the lowest plasma powers in order to ensure precursor

structure retention in the films.

138

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0 10 20 30 40 50 60 70 80 90 100 110

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Figure 62: Increasing film thickness retention with increasing plasma

powers suggests that at higher powers precursor fractionation in plasma yields more densely cross-linked AA films.

Since minimal plasma powers need to be utilized during depositions, an initiator

would need to be added to the precursor mixture to enable depositions at reasonable rates.

tert-butyl peroxide (tBPO) was chosen as the initiator to be used, since it could be easily

broken up to form peroxide radicals without requiring much energy. Depositions were

made at the following conditions: AA flowrate was kept constant at 3 sccm, argon

flowrate was 10 sccm, constant pressure of 250 mTorr was maintained, and plasma was

continuous at 10 W. tBPO flowrate was varied in order to find a condition, at which

polymerization of AA would be enhanced by the addition of the initiator, but would still

yield films with a high concentration of carboxyl groups. It was observed that increasing

the tBPO flowrate increased the degree of cross-linking in the pAA films. Figure 63

139

shows that the thickness retention of the films that were washed in DI water for 10

minutes increased with the addition of tBPO into the reactor.

0

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35

40

45

0 0.05 0.1 0.15 0.2 0.25 0.3tBPO Flowrate (sccm)

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Figure 63: Improvement in film thickness retention with the increase

of tBPO flowrate, which points to tBPO contributing to film cross-linking.

This improvement in film stability in water can be explained by the abundance of

tBPO radicals in the plasma, which then not only initialize the linear polymerization of

AA through the vinyl chemistry, but also react with other available bonds in the growing

pAA linear chains and cross-link these. This hypothesis is supported by the FTIR spectra

of the pAA films with varying amounts of tBPO initiator, presented in Figure 64, which

exhibit a decrease in the intensity in the carbonyl peak, occurring at 1760-1700 cm-1, with

increasing tBPO flowrate, indicating that some of the carbonyl bonds are consumed by

the radicals to initiate cross-linking.

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tBPO - 0 sccm

tBPO - .1 sccm

tBPO - .25 sccm

Increasing tBPO flowrate

Figure 64: Segment of FTIR spectra of pAA films deposited with

varying amounts if tBPO. The decrease in the carbonyl peak intensity (1740-1700 cm-1) with increasing tBPO flowrate pointed to cross-linking occurring in the film due to the presence of tBPO radicals in the plasma.

Protein attachment to a surface is typically performed in an aqueous environment.

This means that the surface, on which the attachment takes place, has to be insoluble in

water. When polymerized, acrylic acid (AA) forms linear carbon chains with carboxylic

acid pendant groups by reacting through the vinyl bond. Due to its chemical structure,

(pAA) is highly hydrophilic and water-soluble. Coating stability in aqueous

environments is a major issue for biocompatibility applications, since it can hinder

protein attachment, and cause further problems with monomer leaching [24-27]. Thus,

pAA films need to be cross-linked to ensure their stability in water. While it was

observed that stability improved with increasing tBPO flowrate, this also caused a

141

decrease in the concentration of the carboxyl groups in the film. On the other hand,

acrylic acid could be copolymerized with ethylene glycol diacrylate (EGDA), which,

thanks to its two acrylic groups, can act as a cross-linker between two linear pAA chains

without causing a loss of carboxyl groups in the film that act as anchoring points for

protein tethering onto the film surface, as pictured in Figure 65.

Figure 65: Illustration of EGDA molecule used as a cross-linker between two polymer chains [62]. For pAA chains R represents a carboxyl group, COOH, while R’ represents a hydrogen molecule.

With this goal in mind pAA-co-EGDA films were deposited with varying

amounts of EGDA. The flowrate of EGDA into the reactor was minimal and difficult to

measure, thus, it was controlled by controlling the temperature of the precursor: between

25 and 75 °C. Depositions were made at following conditions: AA flowrate was kept

constant at 3 sccm, argon flowrate was 10 sccm, constant pressure of 250 mTorr was

maintained, and plasma was continuous at 10 W. tBPO flowrate was kept at a minimal

value of 0.1, which was enough to initialize AA and EGDA precursors at low plasma

powers, but not so high so as to cause reactions between bonds other than vinyl in plasma

and loss of functional groups in the films.

142

It was observed that increasing the EGDA temperature resulted in greater

incorporation of EGDA molecules into the pAA films as evidenced by the carbonyl peak

between 1760 and 1720 cm-1. In AA-only film spectra this peak occurred at about 1710

cm-1, with a shoulder at higher wavenumbers, as can be seen in Figure 66. The peak

associated with the carbonyl group in EGDA was said to appear at around 1730 cm-1,

with a shoulder at lower wavenumbers, according to previous research [61]. Figure 66

shows that the peak at 1735 cm-1, indicating carbonyl groups in EGDA-like structure,

increased in intensity as the temperature of the EGDA precursor increased from 45 to 75

°C.

0.005

0.015

0.025

0.035

0.045

16701680169017001710172017301740175017601770Wavenumbers (cm-1)

Abs

orba

nce

(arb

itrar

y un

its)

no EGDA

EGDA - 45°C

EGDA - 55°C

EGDA - 65°C

EGDA - 75°C

Figure 66: Segment of FTIR spectra of pAA-co-EGDA films deposited

with varying EGDA temperatures. The shift in the carbonyl peak from 1710 cm-1 to 1735 cm-1 points to increasing EGDA incorporation into the film with increasing EGDA precursor temperature.

143

The solubility of the pAA-co-EGDA films tracked with the amount of the EGDA

in the film, and, in turn, with the temperature of the EGDA precursor, which was used to

control the flowrate. Figure 67 shows that the thickness retention of the films that were

washed in DI water for 10 minutes increased from ~7 % for the film with no EGDA at all

to ~ 78 % for the film, which was deposited with EGDA at 75 °C.

0

10

20

30

40

50

60

70

80

90

100

20 25 30 35 40 45 50 55 60 65 70 75 80

Temperature of EGDA (°C)

% T

hick

ness

Ret

entio

n

Figure 67: Increasing film thickness retention with increasing EGDA

precursor temperature points to EGDA acting as a cross-linker in the films.

Once the deposition conditions for pAA-co-EGDA films were optimized the films

had to be combined with V3D3 films that would provide electrical insulation for the

biodevices. This, as well as utilization of carbonyl groups in pAA-co-EGDA films for

protein attachment was done by Stephens [63]. The films were deposited using 2.5 sccm

V3D3 flowrate, 0.5 sccm AA flowrate, 1 sccm tBPO flowrate, 5 sccm Ar flowrate, and

144

EGDA temperature of 58 °C. The deposition was performed with 10 W plasma power.

A V3D3 film was deposited initially with Ar and tBPO, after which V3D3 flow was

turned off and AA and EGDA flows were turned on, so as to provide a film with separate

sections for insulation and biofunctionalization. A film with total thickness of 781 nm

was coated onto a glass slide.

The FTIR spectrum of the film, presented in Figure 68, showed the presence of

siloxane peaks between 1300 and 1000 cm-1 and a carbonyl peak at 1740-1700 cm-1,

proving that the film was made up of an electrically insulating V3D3 layer and a

functional pAA-co-EGDA layer. The hydroxyl peak occurring between 3600 and 3100

cm-1 could also be attributed to the acrylic acid in the film.

-0.01

0

0.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08

0.09

0.1

400900140019002400290034003900

Wavenumbers (cm-1)

Abs

orba

nce

(arb

itrar

y un

its)

C=O

CHx

-OH

-Si-O-Si-

Figure 68: FTIR spectrum of V3D3/pAA-co-EGDA film.

145

The portion of the FTIR spectrum between 1300 and 1000 cm-1 attributed to

V3D3 in the film, presented in Figure 69, can be separated into several different peaks.

The peak occurring between 1100 and 1000 cm-1 is formed of peaks at ~1015, ~1035, and

~1070 cm-1. The 1035 and 1170 cm-1 peaks can be ascribed to long- and short-chain

siloxane structures with terminal methyl or hydroxyl groups, present in the film due to

precursor structure loss of V3D3 in plasma. The peak occurring at 1070 cm-1 can be

attributed to long siloxane chains in the film, also forming as a result of precursor

fractionation in plasma. The peak of greatest intensity at ~1015 cm-1 is typically assigned

to cyclotrisiloxane structures [41, 42], and suggests that most of the siloxane moieties in

the film are still in cyclic form. Another peak of interest in the siloxane range of the

spectrum is that occurring at 1280-1250 cm-1, which is attributed to oxygen substitution

on silicon atoms, and shifts towards higher wavenumbers with increasing number of

oxygen atoms. The exact position of the peak in the region points out the degree of

oxidation of the silicon atoms: a peak centered at ~1250 cm-1 suggests mono-substituted

groups, at ~1260 cm-1, di-substituted groups, and at ~1270 cm-1, tri-substituted groups.

The absorbance from quad-substituted groups does not appear in this region but rather

between 1200 and 1000 cm-1, since there are no silicon-organic bonds in these groups.

As can be seen in Figure 69, the peak in the V3D3/pAA-co-EGDA film spectrum is

centered close to 1260 cm-1, suggesting that the greatest contribution is from di-

substituted groups. This fits well with the molecule structure of V3D3 precursor where

each silicon atoms is bonded to two oxygen atoms and two organic groups. The oxygen

substitution peak has a shoulder at 1270 cm-1, due to tri-substituted silicon groups, which

points to V3D3 fragmentation in plasma, and some densification of the film structure.

146

The carbonyl range in the FTIR spectrum presented in Figure 70 could be

deconvoluted into two peaks: 1710 and 1735 cm-1, which can be attributed to AA and

EGDA structures in the film. As previously mentioned, incorporation of EGDA into the

film is important, as it ensures film stability in aqueous media.

0

0.005

0.01

0.015

0.02

0.025

0.03

0.035

9509751000102510501075110011251150117512001225125012751300

Wavenumbers (cm-1)

Abs

orba

nce

(arb

itrar

y un

its)

Figure 69: Siloxane range in the FTIR spectrum of V3D3/pAA-co-

EGDA film.

147

0.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08

0.09

0.1

1680169017001710172017301740175017601770

Wavenumbers (cm-1)

Abs

orba

nce

(arb

itrar

y un

its)

Figure 70: Carbonyl peak in the FTIR spectrum of V3D3/pAA-co-

EGDA film.

The film stability in aqueous environment was tested by soaking it in DI water for

1 hr, after which the film thickness retention was approximately 95 %. This was thanks

to good cross-linking in the film, due to copolymerization of AA and EGDA, as

evidenced by the presence of both AA- and EGDA-specific carbonyl peaks in the FTIR

spectrum in Figure 70, as well as likely grafting between the V3D3 film and the pAA-co-

EGDA top layer. The stability of the film in water proved that it was suitable for further

protein attachment experiments.

Protein attachment was performed using the “PolyLink Protein Coupling Kit for

COOH Microspheres” from Bangs Laboratories, Inc. The glass slide with the V3D3/AA

coating was immersed into the wash buffer [10 mM tris(hydroxymethyl)aminomethane

(tris), pH 8.0; 0.05 % Bovine Serum Albumin; 0.05 % Proclin © 300] for 15 minutes. A

148

carbodiimide solution was prepared using a ratio of 15 mg carbodiimide to 1.5 mL

coupling buffer [50 mM 2-(N-morpholino)ethanesulphonic acid (MES), pH 5.2; 0.05

Proclin © 300 (antimicrobial preservative for enzymes)]. The glass slide was immersed

into the carbodiimide solution, allowed to react for 4 hours, and then rinsed with the wash

buffer. It was then immersed into the coupling buffer and fluorescein isothiocyanate

(FITC) anti-mouse immunoglobulin (IgG) fluorescently labeled protein was added in the

volumetric ratio of 2 µL protein to 1 mL buffer. The glass slide was again rinsed with the

wash buffer and examined under fluorescent microscope.

Fluorescent microscopy showed the FITC anti-mouse IgG protein attached to the

surface of the V3D3/AA coated glass slide, indicated by the red glow in Figure 71. This

confirmed that it is possible to combine electrically insulating and biocompatible and

biofunctional surfaces into one film using CVD techniques.

149

Figure 71: Fluorescent micrograph of FITC anti-mouse IgG protein

tethered to V3D3/pAA film surface.

150

5. Conclusions

The goal of this investigation was to obtain a thin film with lowest possible

dielectric constant. The refractive index, which was comparable to the dielectric constant

of the films, was shown to be dependent on the polizability and the density of the film.

Organosiloxane precursors were used to provide films with low polarizability. The

density of the films was to be reduced by incorporating pores into the film. V3D3 was

used as the monomer of choice with intrinsic porosity. Induced porosity was obtained by

adding linear co-monomers, such as VDEMS, S1, S3, and S5, to the V3D3 film. The

good insulating properties of organosiloxane films, together with their biocompatibility

shown in literature, made them good candidates as coatings for bioimplantable devices.

In addition to obtaining a film with low dielectric constant, this study also explored

depositions of V3D3 film with a top layer of pAA that could be functionalized with

biomolecules, as well as the stability of these layered films in aqueous environment. Due

to equipment availability refractive index was measured in lieu of dielectric constant,

since the two quantities were related.

PECVD and iCVD techniques were used for deposition of V3D3 and V3D3-

spacer films. PECVD was shown to be too aggressive as a deposition technique that

resulted in precursor structure loss in the films even when mildest deposition conditions

were used. The lowest refractive index obtained for a V3D3 film deposited at a low

substrate temperature and low power was 1.482 ± 0.001. This value was reduced

substantially once the film was annealed in a nitrogen atmosphere for 90 minutes at 400

°C: the post-anneal refractive index was 1.422 ± 0.001. By contrast the mildest iCVD

151

conditions yielded a V3D3 film with significant retention of cyclic structure, as

evidenced by FTIR spectra and peak deconvolution, and with a refractive index of 1.454

± 0.002, much lower than that of the PECVD film. This showed that iCVD was a more

suitable milder alternative for depositing films from precursors with delicate

functionality.

Deposition of a V3D3 film at a high substrate temperature of 200 °C using the

PECVD technique resulted in a film that did not retain much of the V3D3 cyclic

structure, but the high substrate temperature had the same effect on the growing film as

annealing that was done after deposition. It allowed the film to “heal” as it was being

deposited, minimizing the number of dangling bonds, and resulting in a more regular

structure, and in a lower refractive index: 1.425 ± 0.005.

Additionally, the composition, properties and thickness changed minimally upon

annealing for the PECVD films deposited at high substrate temperature, while the

ambient substrate temperature films exhibited not only significant reduction in refractive

index values, but also a major change in film structure, seen spectrally, which was

coupled with thickness losses as high as 20 %.

VDEMS had an interesting molecular structure and was thought to make a good

potential spacer molecule for films with induced porosity. The VDEMS-only film with

best optical properties was deposited at 100 W plasma power, 1 sccm flowrate, and

ambient stage temperature. While it had a refractive index of 1.470 as-deposited, it

dropped down to 1.411 post-anneal, with a 7 % thickness loss in the film. The VDEMS

films deposited at high substrate temperatures had fairly low refractive indices as-deposit,

~1.430, but these remained unchanged post-anneal, thus, showing that there was a trade-

152

off between thermal stability and optical properties. Deposition of a V3D3-VDEMS film

at high substrate temperature, at conditions that had yielded successful V3D3 and

VDEMS films, produced a film with a dielectric constant of 1.432 ± 0.002, which did not

change significantly upon annealing: 1.422 ± 0.005, leaving more to be desired from

V3D3-spacer films.

V3D3 was copolymerized with a series of linear siloxane molecules, which were

intended to act as spacers to further lower the refractive index of the films. The

depositions using the PECVD technique yielded films, refractive indices of which were

higher than those of V3D3-only films deposited at the same conditions, both as-deposit

and post-anneal. For example, one of the most successful PECVD V3D3-spacer films

was the V3D3-S3 film deposited at 200 °C and 10 W, which had a post-anneal refractive

index of 1.422 ± 0.001, while the refractive index of a post-anneal V3D3 film deposited

at the same conditions was 1.420 ±0.002. iCVD yielded V3D3-spacer films that had

lower refractive indices than their PECVD counterparts deposited at low substrate

temperature. However, since none of the iCVD films were annealed, the true extent of

improvement of changing the deposition technique was difficult to illustrate.

The fact that refractive index values of V3D3-spacer films were higher than those

of V3D3-only films deposited at the same conditions suggested that the linear siloxane

molecules were not acting as spacer molecules as had been hypothesized. It was likely

that the spacers polymerized through the vinyl bonds only, similar to acrylates, so that the

linear part of the molecule acted as a pendant group. This would explain and the increase

in the refractive index with addition of the spacers to V3D3 films as well as the increase

in intensity of the short-chain/cage structure peaks. Additionally, in PECVD the

153

abundance of energy could have resulted in precursor fragmentation leading to film

densification, as opposed to additional pore formation.

PM/Psat ratio for the monomer was proven to be a crucial variable in iCVD by Lau

and Gleason [52, 53], particularly where co-polymerization was concerned, since it

quantified the amount of monomer adsorbed onto the substrate, which, in turn, controlled

the deposition rate. The deposition rate was shown to increase with increasing PM/Psat

value. It was shown that a V3D3-S5 film that deposited with equal PM/Psat ratios (0.42)

for both monomers had a refractive index of 1.453 ± 0.010, whereas the film deposited

with PM/Psat value of 0.15 for V3D3 and 0.23 for S5 had a refractive index of 1.461 ±

0.002 respectively. It was difficult to judge on the effect of the PM/Psat ratio on the

optical properties of the films, as nearly all refractive index values were within the error

margin.

V3D3 films were investigated as base layers for biofunctionalized films.

Biofunctionalization would be achieved by protein attachment to carboxylic acid groups

in the top pAA layer. In order to serve as a successful biofunctionalized coating the pAA

film had to not only retain a significant number of carboxylic acid groups on the surface,

but also be stable in aqueous environment. Thus, pAA films were co-polymerized with

EGDA that acted as a cross-linker. It was shown that increasing plasma power improved

the films’ stability in water, but resulted in the decrease of the carbonyl peak in their

FTIR spectra, suggesting that cross-linking took place at the expense of the carboxyl acid

functionality. It was also shown that increasing the flowrate of tBPO initiator resulted in

more stable films in aqueous environments, but this also resulted in simultaneous drop off

in carbonyl peak intensity in the FTIR spectra. This was hypothesized to be due to the

154

fact that the abundance of tBPO radicals in the plasma not only initialized the linear

polymerization of AA through the vinyl chemistry but also reacted with other available

bonds in the growing pAA linear chains and cross-linked these, thus, consuming

carboxylic acid functional groups. Co-polymerization with EGDA showed to have the

desirable effect on film stability without depleting the films of their important

functionality. It was also shown that increasing the EGDA flowrate into the reactor

(done by increasing the temperature of the precursor in the jar before it was allowed to

flow into the reactor) resulted in more stable films.

Based on some additional work done by Stephens [63] within the framework of

the same project it was observed that a combination V3D3/pAA-co-EGDA film had good

thickness retention, ~95 %, when soaked in DI water for an hour. This was thanks to

good cross-linking in the film, due to copolymerization of AA and EGDA, as well as due

to likely grafting between the V3D3 film and the AA/EGDA top layer.

Protein tethering was performed on the same combination V3D3/pAA-co-EGDA

film using FITC anti-mouse IgG protein, which was shown to be attached to the surface

by fluorescent microscopy proving that was is possible to combine electrically insulating

and biocompatible and biofunctional surfaces into one film using CVD techniques.

155

6. Future Work

Some of the recommended work includes annealing iCVD films to have a better

frame of reference for the degree of cross-linking in these films. This will also allow for

the comparison of post-anneal refractive index values between PECVD and iCVD films,

since the lowest refractive index values were obtained for annealed PECVD films. Most

importantly, it will allow to relate cyclic structure retention to cross-link formation in the

films.

More efforts need to be concentrated understanding the role of the PM/Psat ratio in

iCVD co-polymerizations and deposition conditions need to be optimized for V3D3-S5

monomer couple. Other spacer monomers need to be research, which have saturated

vapor pressures close in value to that of V3D3, so that the PM/Psat ratio value can be the

same for both monomers during deposition. Additionally, since hypothesis about the

linear siloxane chain monomers with end-capping vinyl groups polymerizing in such a

way as to become spacers between V3D3 rings may be false, other potential spacer

moieties need to be explored. A good candidate would be a similar linear siloxane

molecule with only a single vinyl, or with no vinyl bonds at all.

Other recommended future work includes depositing V3D3/pAA-co-EGDA films

using iCVD, since it would allow for a greater extent of retention of delicate carboxylic

acid functional groups and also would yield V3D3 base film with better insulating

properties. Also, the relationship between structure retention and degree of cross-linking

in the films would need to be studied since the biofunctionalization step takes place in

aqueous environment.

156

7. Nomenclature

A Absorbance

[C] Concentration

[C]0 Initial concentration

[C]t Concentration at time t

c Speed of light

E Young’s modulus

E1, E2 Energy states

h Planck’s constant

I Intensity of the laser source

I0 Partial beam intensity

Ion, Ioff Peak intensities of molecule fragments measured with and without plasma

k Relative dielectric constant

k0, k1, k2 Cauchy extinction coefficients

kobs Observed reaction rate

kr Dielectric constant of the porous material

kl Dielectric constant of material inside pores

ks Dielectric constant of film skeleton

L Length of the metal wire

l Beam path length

N Number of molecules per unit volume

Ñ Complex index of refraction

157

n Index of refraction

n0, n1, n2 Cauchy parameters

P Porosity of the material

P Metal line pitch (equation 4)

PM Monomer gas pressure

Psat Saturated vapor pressure of the monomer

RC Resistive capacitive delay

T Thickness of metal line

t Time

x Ratio of squares of the longitudinal frequency to the transverse frequency

α Total polarizability

ε Permittivity

ε Molar absorptivity (equation 9)

ε0 Vacuum permittivity

η1, η2 Fractions of materials 1 and 2 in the system

λ Wavelength

λ Retardation (equation 13)

ν Frequency of light

˜ ν Wavenumber

ρ Metal resistivity

158

8. Abbreviations

AA Acrylic acid

CPO Cyclopentene oxide

CVD Chemical vapor deposition

D3 Hexamethylcyclotrisiloxane

D4 Octomethylcyclotetrasiloxane

D5 Decamethylcyclopentasiloxane

DEMS Methyldiethoxysilane

DMDMOS Dimethyldimethoxysilane

DVDMDSO Divinyldimethyldisiloxane

EGDA Ethylene glycol diacrylate

FTIR Fourier transform infrared (spectroscopy)

H4D4 Tetramethylcyclotetrasiloxane

HFCVD Hot filament chemical vapor deposition

HMDSO Hexamethyldisiloxane

iCVD Initiated chemical vapor deposition

MEMS Microelectromechanical systems

MMA Methyl methacrylate

MS Mass spectrometry

MTMOS Methyltrimethoxysilane

OES Optical emission spectroscopy

PC12 Cell line derived from a pheochtomocytoma of the rat adrenal medulla

159

PECVD Plasma enhanced chemical vapor deposition

SSQ Silsesquioxane

tBPO tert-Butyl peroxide

TMMOS Trimethylmethoxysilane

TMS Tetramethylsilane

TVMS Trivinylmethylsilane

ULSI Ultra large scale integrated (devices)

V3D3 Trivinyltrimethylcyclotrisiloxane

V4D4 Tetravinyltetramethylcyclotetrasiloxane

V5D5 Pentavinylpentamethylcyclopentasiloxane

VDEMS Vinylmethyldiethoxysilane

VTMS Vinyltrimethylsilane

XPS X-ray photoemission spectroscopy

160

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