ADHERENCE/DIFFUSION BARRIER LAYERS FOR COPPER METALLIZATION:

AMORPHOUS CARBON:SILICON POLYMERIZED FILMS

Merry Pritchett, M.S.

Dissertation Prepared for the Degree of

DOCTOR OF PHILOSOPHY

UNIVERSITY OF NORTH TEXAS

May 2004

APPROVED: Jeffry A. Kelber, Major Professor Robert Desiderato, Committee Member Teresa Golden, Committee Member Rick Reidy, Committee Member Ruthanne D. Thomas, Chair of the Department of

Chemistry Sandra L. Terrell, Interim Dean of the Robert B.

Toulouse School of Graduate Studies

Pritchett, Merry, Adherence/Diffusion Barrier Layers for Copper Metallization:

Amorphous Carbon:Silicon Polymerized Films, Doctor of Philosophy (Analytical

Chemistry), May 2004, 102 pp., 4 tables, 35 illustrations, reference list, 127 titles.

Semiconductor circuitry feature miniaturization continues in response to Moore’s

Law pushing the limits of aluminum and forcing the transition to Cu due to its lower

resistivity and electromigration. Copper diffuses into silicon dioxide under thermal and

electrical stresses, requiring the use of barriers to inhibit diffusion, adding to the

insulator thickness and delay time, or replacement of SiO2 with new insulator materials

that can inhibit diffusion while enabling Cu wetting. This study proposes modified

amorphous silicon carbon hydrogen (a-Si:C:H) films as possible diffusion barriers and

replacements for SiO2 between metal levels, interlevel dielectric (ILD), or between metal

lines (IMD), based upon the diffusion inhibition of previous a-Si:C:H species expected

lower dielectric constants, acceptable thermal conductivity. Vinyltrimethylsilane (VTMS)

precursor was condensed on a titanium substrate at 90 K and bombarded with electron

beams to induce crosslinking and form polymerized a-Si:C:H films. Modifications of the

films with hydroxyl and nitrogen was accomplished by dosing the condensed VTMS with

water or ammonia before electron bombardment producing a-Si:C:H/OH and a-Si:C:H/N

and a-Si:C:H/OH/N polymerized films in expectation of developing films that would

inhibit copper diffusion and promote Cu adherence, wetting, on the film surface. X-ray

Photoelectron Spectroscopy was used to characterize Cu metallization of these a-

Si:C:H films. XPS revealed substantial Cu wetting of a-Si:C:H/OH and a-Si:C:H/OH/N

films and some wetting of a-Si:C:H/N films, and similar Cu diffusion inhibition to 800 K

by all of the a-:S:C:H films. These findings suggest the possible use of a-Si:C:H films

as ILD and IMD materials, with the possibility of further tailoring a-Si:C:H films to meet

future device requirements.

ii

Copyright 2003

by

Merry Pritchett

iii

ACKNOWLEDGEMENTS

I would like to thank my University of North Texas thesis committee, Dr. Teresa

Golden, Dr. Rick Reidy, Dr. Robert Desiderato, and Dr. Moon Kim from the University of

Texas – Dallas. I appreciate your insights, comments and suggestions.

I especially would like to express my appreciation to my advisor, Dr. Jeff Kelber,

for being such a “good guy” – working with my teaching schedule so that I could finish

my master’s degree and for giving me the opportunity to join his group and do really

interesting doctoral research. I’d also like to thank him for encouraging me to utilize

resources from other groups, enabling me to pursue a pharmaceutical internship, and

providing access to financial support from the Semiconductor Research Corporation

(SRC) -TI/SRC Customized Research, and Robert Welch Research Grants.

Much appreciation also goes to those who provided resources for me to extend

my research: Rick Reidy, for the use of the Materials Synthesis and Processing FTIR

instrument, Brian Gorman, for his FTIR expertise, Bob Wallace and Prakaipetch

Punchaipetch, for access to the Laboratory for the Electronic Materials and Devices.

I am truly grateful to Dr. Ed Dorsey for creating an internship at Alcon

Laboratories, Inc. where there was none, and Brett Scott, Leon Donesky, Ivy Badger

and Phillip Ritter, and Dr. Danny Dunn, for providing a year of invaluable, enjoyable,

learning experiences.

Immeasurable thanks go to my parents for memorable summer trips to national

parks and museums that cultivated in me a lasting appreciation for science and for their

constant encouragement and support during my studies.

iv

This thesis is dedicated to the memory of my grandmother, who constantly

encouraged me to “Hang in there”, to my father, who hung in there for me, and to God,

the creator of atoms, interactions and materials – Jehovah-jireh.

v

TABLE OF CONTENTS

Page ACKNOWLEDGMENTS ......................................................................................... iii LIST OF TABLES ................................................................................................... vii LIST OF ILLUSTRATIONS ..................................................................................... viii Chapter

1. INTRODUCTION .................................................................................... 1 1.1 Interconnect Structure and Resistance-Capacitance Delay ...... 1 1.2 Metallization of Interconnects ................................................... 5 1.3 Diffusion Barriers ...................................................................... 6 1.4 Dielectrics ................................................................................. 9 1.5 Adhesion ................................................................................... 11 1.6 Techniques and Instrumentation .............................................. 15

1.6.1 X-ray Photoelectron Spectroscopy (XPS) ...................... 15 1.6.2 Physical Vapor Deposition (PVD) .................................. 22 1.6.3 Precursor Deposition ...................................................... 25 1.6.4 XPS, PVD and Precursor Doser Apparatus .................... 26

1.7 Chapter References ................................................................. 28

2. HYDROXYLATED AMORPHOUS POLYMERIZED CARBON:SILICON FILMS - ADHERENCE /DIFFUSION BARRIERS FOR COPPER ........ 35

2.1 Motivation for Si:C:H Based Barrier Layer ................................ 35 2.2 Experiment ............................................................................... 38 2.3 Results ...................................................................................... 41

2.3.1. Amorphous-Si:C:H Film Characterization ....................... 41 2.3.2. Copper Deposition by Physical Vapor Deposition (PVD) 44 2.3.3. Annealing Effects ............................................................ 48

2.4 Discussion ................................................................................. 50 2.5 Conclusions .............................................................................. 54 2.6 Chapter References .................................................................. 56

vi

3. NITROGEN MODIFIED AMORPHOUS POLYMERIZED CARBON:SILICON FILMS - ADHERENCE /DIFFUSION BARRIERS FOR COPPER ........ 60

3.1 Motivation for Nitrogen Modification of a-Si:C:H Films ............. 60 3.2 Experiment ............................................................................... 63 3.3 Results ...................................................................................... 66

3.3.1. Nitrogen Modified a-Si:C:H Film Characterization .......... 66 3.3.2. Copper Deposition by Physical Vapor Deposition (PVD) 70 3.3.3. Annealing Effects ............................................................ 77

3.4 Discussion ................................................................................. 79 3.5 Conclusions .............................................................................. 82 3.6 Chapter References .................................................................. 84

4. RESEARCH SUMMARY ........................................................................ 88 4.1 Chapter References .................................................................. 91

REFERENCES........................................................................................................ 92

vii

LIST OF TABLES Table Page 1.1. Aluminum and Copper properties .............................................................. 5 1.2. Comparison of thermal mechanical properties of SOD and SiOC films .... 8 1.3. Dielectric constants of common dielectric materials................................... 10 2.1 Relative surface concentrations for Si, Cl and C on Al at different ............ CTCS coverages at room temperature ...................................................... 39 3.1. Binding Energies of a-Si:C:H Film Components (eV)................................. 67

viii

LIST OF ILLUSTRATIONS Figure Page 1.1. Tungsten via plug.................................................................................... 1 1.2. Comparison of intrinsic gate delay and interconnect (RC) delay as a function of feature size. ................................................................... 2 1.3. Microprocessor clock speed for different metals and dielectrics as a function of feature size. ................................................................... 3 1.4. Hillock and void formation ....................................................................... 4

1.5 Interconnect metallization structure......................................................... 6 1.6. Wetting described by contact angle and temperature ............................. 12 1.7. Horizontal approximation of the 180º Peel Test ...................................... 14 1.8. XPS electron emission process............................................................... 16 1.9. Nitrogen adsorbed on Ni(100) surface. ................................................... 17 1.10. Comparison of angle resolved and normal emission geometry............... 18

1.11. XPS Instrumentation. Angle resolved, θ = 15°, Si 2p spectrum ............. enhances the surface oxide species as compared to Normal Si 2p ....... spectrum, θ = 90°.. .................................................................................. 19 1.12. XPS Instrumentation. .............................................................................. 20 1.13. Physical Vapor Deposition process......................................................... 22 1.14 Three methods of metal film growth. (a) Frank van de Merwe (b) ........ Stranski - Krastanov (c) Volmer Weber................................................... 24 1.15. Uptake curves of Metal/Substrate versus metal deposition time depicting two

different growth modes: a) conformal growth b) 3-D island growth ..... 25 1.16. XPS and CVD and PVD system26 2.1. XPS spectra of unmodified a-Si:C:H film................................................. 39 2.2. XPS spectra of hydroxyl-modified a-Si:C:H film ...................................... 40

ix

2.3. XPS spectra of Cu deposition on unmodified a-Si:C:H film .................... 43 2.4. XPS spectra of Cu deposition on OH-modified a-Si:C:H film ................. 43 2.5. Cu(2p3/2)/Si(2p) uptake curve.................................................................. 44 2.6. Cu(L3VV) Auger peak XPS spectra of hydroxyl modified a-Si:C:H film ... during Cu deposition: ............................................................................ 45 2.7. Hydroxyl modified a-Si:C:H film intensity ratios versus anneal ............... temperature............................................................................................. 47 2.8. Cu Auger parameter indicates both Cu(I) and Cu(0) species are ........... deposited................................................................................................. 50 3.1. Unmodified a-Si:C:H film showing single O(1s), C(1s) and Si(2p) species ................................................................................................... 65 3.2. Comparison of XPS spectra (a) nitrogen modified a-SI:C:H film and (b) nitrogen-hydroxyl modified a-Si:C:H film. .............................................. 66 3.3 XPS Cu(2p) spectra showing absence of Cu(II) shake-up (satellite ) ..... peaks ...................................................................................................... 69 3.4 XPS spectra of copper deposition on nitrogen/hydroxyl modified .......... a-Si:C:H film ............................................................................................ 70 3.5. XPS Cu(2p)/Si(2p) Intensity ratio - Copper uptake curve: a) a-Si:C:H/N film b) a-Si:C:H film ................................................................................ 71 3.6 Cu(2p 3/2)/Si(2p) uptake curve indicating a change in deposition mode .. as copper is deposited on a-Si:C:H/N films............................................. 72 3.7. Cu(2p 3/2)/Si(2p) uptake curve indicating a change in deposition ............ mode as copper is deposited on a-Si:C:H/OH/N films............................. 72 3.8. Copper deposited on a-Si:C:H/N films and a-Si:C:H films show only .... Cu(0) species shifting to higher kinetic energy due to increase in cluster size and final state shielding ................................................................... 69 3.9. Copper deposited on a-Si:C:H/OH/N films shows Cu(I) and Cu(0) ........ species simultaneously, producing the broad peak shape. Copper ...... deposited on a-Si:C:H films show only Cu(0) species, indicated by ....... narrow peak shape that shifts to higher kinetic energy due to increase . in cluster size and final state shielding .................................................... 70

x

3.10. Copper diffusion is signaled by the rapid increase in the Cu(3p)/Cu(2p) Ratio........................................................................................................ 72

3.11. Copper Auger intensity decreases with anneal from 300 K –1000 K, ..... and the peak becomes sharper as it becomes more metallic.................. 73

1

CHAPTER 1

INTRODUCTION

1.1 Interconnect Structure and Resistance – Capacitance Delay

Interconnects are metal circuits that carry electron signals between transistors

and are basic to semiconductor technology. Multilayer structures are used to enable

smaller areas to incorporate more circuitry. Metal lines consisting of vias delineated by

Fig. 1.1 Tungsten via plug [1] (Used with permission, Thompson).

dielectric interlayers, alternate on a tungsten or silicon substrate., ending with a polymer

layer, capping layer, to keep corrosion and extraneous electrons out [2]. As

technological advancements enable semiconductor chips to get smaller, the smaller

feature sizes of chip components require much smaller integrated circuits and metal

2

lines connecting those circuits. The size of metal lines in interconnects significantly

affects the speed of electron flow and is based upon the metal’s resistance, R, cross

sectional area, A, length, L, and resistivity, ρ:

R = ρ L A

Therefore decreasing line length and area increases resistance, due to line width

constraints. The type of dielectric material largely determines the capacitance, C, given

by the dielectric constant, K, area, A, and thickness, t,:

C = K * (A/t) Decreasing feature sizes will decrease the A/t ratio, requiring lower dielectric constant

materials to achieve lower capacitance.

Intrinsic material properties significantly affect the conductivity of electrons as

feature sizes decrease. Two parameters that affect the speed an electrical signal are

Fig. 1.2 Comparison of intrinsic gate delay and interconnect (RC) delay as a function of feature size [3] (Used with permission, Barnes).

3

carried are transistor gate delay (the switching time of an individual transistor) and

resistance-capacitance (R-C) delay; the signal propagation time between transistors.

As shown in Figure 1.2, interconnect delay (R-C delay) increases as feature sizes

decrease, whereas transistor gate delay decreases. R-C delay is dependent on the

metal properties, feature sizes and properties of the dielectric [3]:

RC delay = 2ρε(4L²/P²+L²/T²)

R = metal wire resistance L = line length C = interlevel dielectric capacitance P = metal pitch ρ = metal resistivity T = metal thickness ε = ILD permittivity, or dielectric constant

The continued shrinkage of interconnect feature size has greatly increased the R-C

delay relative to gate delay (Fig. 1.2) and has motivated the switch from aluminum to

Fig. 1.3 Microprocessor clock speed for different metals and dielectrics as a function of feature size [3] (Used with permission, Barnes).

Cu-Low κ Al-Low κ

Cu-SiO2

Al-SiO2

Technology Min. Feature (nm)

250 180 150 30 100 70 50

Clo

ck F

req.

(M

Hz)

3200

2800

2400

2000

1600

1200

800

400

0

4

copper interconnects (lower R) and from SiO2 to lower dielectric constant insulators

(lower C) requiring the adoption of new materials and processes. In the transition from

aluminum to copper, and resultant adoption of new materials, the speed of processors

is related to R-C delay more than transistor gate delay (Fig. 1.3).

Aluminum has been the metal of choice on dielectric silicon dioxide due to an

interfacial passivation reaction that enables aluminum to adhere to SiO2 without

diffusing into the bulk. Increases in resistance and signal processing time due to

feature reduction, necessitate a metal with a lower intrinsic resistance than aluminum.

In addition, aluminum atoms have a tendency to migrate along with the flow of

electrons, electromigration, leaving voids at the starting point and piling up at the

Fig. 1.4 Hillock and void formation [1] (Used with permission, Thompson).

5

end point, creating hillocks (Figure 1.4) [4]. Both of these problems are remedied by

changing to copper. Copper is more conductive, with greater resistance to

electromigration and a higher melting point than aluminum as shown in Table 1.1.

The replacement of aluminum by copper in interconnects poses several

problems. First Cu does not form a passivation layer, therefore presenting a possible

oxidation problem, and it poorly adheres to most of the most commonly used dielectric

material, SiO2. In addition, copper readily diffuses into SiO2 [5] under conditions of

Table 1.1 Aluminum and Copper Properties (Used with permission, NIST).

N.A. Lange, Handbook of Chemistry pp. 100-107 (pp. 854-861) (Handbook Publishers, Sandusky, Ohio, 1956), *Thermometric fixed point; E.A. Brandes, Smithells Metals Reference Book, 6th ed., p. 16-2 (Butterworths, London, 1983)

combined bias and thermal stress (BTS) due to the lack of formation of an interfacial

passivation layer. In addition, copper does not wet SiO2. Therefore, copper requires

Metal At. Wt.

Density (20° C) (g/cm3)

Melt. Point (° C)

Heat capacity

(25° C) J/(kg.K)

Thermalexp. coef.

(20° C) (10-6 K-1)

Electrical Resistivity

(20° C) (10-6 W .cm)

Thermal conductivity

(W/(m.K))

Aluminum 26.98 2.702 660.46* 902.1 23.03 2.62 238

Copper 63.54 8.92 1084.9* 385.2 16.6 1.69 416

Gold 196.97 19.3 1064.4* 127.5 14.2 2.4 311

Silver 107.87 10.5 961.93* 236.3 18.9 1.62 417

6

the use of a barrier material that promotes copper adhesion and conformal growth

(wetting) as well as the inhibition of its diffusion into the dielectric.

1.3 Diffusion Barriers

In the past a number of different species have been explored as barrier layers,

e.g. TiN/Ti [6, 7] and TaN/Ta [8], or modifications of these e.g. W/TiN [9], TiSiN [10].

The problem with Ta barrier layers is that they can be contaminated with small amounts

of carbon and oxygen, inhibiting copper wetting and adhesion [11]. The formation of Ta

oxide at grain boundaries may result in the loss of Cu film continuity resulting in the

Fig. 1.5 Interconnect metallization structure [1] (Used with permission, Thompson).

7

formation of Cu silicide [12]. Therefore, Ta and TaN dewetting is foreseeable on the

proposed low-k dielectrics without stringent control of background contaminants during

copper deposition [13].

Other proposed barrier layers include Co(W,P) thin films [14], amorphous silicon

carbide (a-Si-C:H) films [15], and silicon oxynitride (SiOxNy) films [16]. Organic films

would be less reactive than Ta, but generally exhibit only weak bonding to Cu and

inhibition of Cu diffusion under Bias-Thermal-Stress conditions raises concerns about

film porosity and thermal and electrical stability. Additionally, the metal barrier and

dielectric substrate may have very different coefficients of thermal expansion, leading to

stress-induced delamination or cracking during thermal analysis. Non-metallic barriers

would have a significant advantage with respect to metal barriers, since their

coefficients of expansion would be more closely matched.

Amorphous silicon carbide barrier layers are formed by dissociating

trimethylsilane (TMS) gas by plasma enhanced chemical vapor deposition (PECVD), at

temperatures from 300 to 400 °C [17]. TMS films deposited on high conductivity silicon,

with gate electrodes of Cu, metal-insulator-semiconductor (MIS) capacitors, were

investigated by capacitance-voltage (C-V) and current-voltage (I-V) measurements.

These films resisted copper diffusion at a voltage of 2 MV for almost 105 seconds,

indicating a higher resistance than plasma silicon nitride, and after bias-temperature-

stress analysis, showed no evidence of copper diffusion [17, 18]. Trimethylsilane

(TMS) low-k oxide films, containing both Si-CH-Si and Si-O-Si bond linkages, are low-

density oxide materials with leakage currents and breakdown strengths similar to SiO2.

8

Stress-temperature analysis that show no apparent change in properties, with a-Si:C:H

used as the barrier layer, suggest their use as interlayer dielectrics [17].

SiOC films have been suggested as low-k barrier films due to their stability to

humidity and integration, superior hardness as compared to spin on films and thermal

characteristics similar to SiO2 as shown in Table 1.3 [19]. In addition, they do not form

pinholes or require solvent removal, they fill gaps well, and their thickness is easily

controlled. Such films, deposited by chemical vapor deposition (CVD), are compatible

with current processing equipment. In contrast, spin on dielectrics (SOD) can be

specifically designed, and have well-known structures such as Dow Corning’s FOx®

Flowable Oxides, hydrogen silsesquioxane (HSQ) [20], and Honeywell’s

NANOGLASS® E Extra Low-k Spin-On Dielectric [21]. However, they require solvent

removal, produce many pinholes, must be cured, often poorly control thickness and

SiO2 SiOC Organic SOD

Hardness (GPa) 9.0 2.0 0.3

Elastic modulus (Gpa) 80.0 10.0 2.5

Thermal conductivity (W/mk) 0.90 0.35 0.20

Thermal expansion CTE (ppm) 0.50 12 62

Table 1.2 Comparison of thermal mechanical properties of SOD and SiOC films [19] (Used with permission, Semiconductor Fabtech).

9

inadequately fill gaps [22]. Considering the mechanical and thermal advantages, as

compared to the technological requirements, SiOC films appear to be much better

choices for use as barrier layers than SOD films.

1.4 Dielectrics

Dielectrics, in a passive mode, insulate and isolate conductors from adjacent

conductors and underlying silicon. Ideally, they should act to hermetically seal

conductors and devices from impurities such as water or corrosive ions that lead to

device instabilities. Passivating films should have few pinholes and keep out moisture.

Barriers to silicon should keep out alkaline ions, or scavenge them. Two reliability

measures for insulators are current leakage and breakdown.

Silicon dioxide is the most commonly used dielectric for aluminum metallization.

However, decreasing feature sizes necessitates dielectrics with lower resistivity, to

counteract “cross-talk” between closely spaced metal lines resulting from capacitive

coupling. Changing the dielectric material from silicon dioxide has initiated the

development of inter-layer and intra-layer (gap-fill) dielectrics of various organic

polymers, e.g. polystyrene or polyimide as (Table 1.3) [23]. Polyimides have the

undesirable ability to incorporate copper in the films, leading to Cu diffusion under BTS

conditions.

Other interlevel dielectrics (ILDs), dielectrics separating electrical conductors,

include the TMS low-k oxide film (silicon oxynitride), because of its similarity to silicon

dioxide, and other low-k films such as Dow Chemical’s SiLK™ low-k spin on dielectric

10

films [24], that do not require adhesion promoters. Low-k methylsilsesquioxane (MSQ)-

based α-SiOC:H films, deposited by plasma enhanced chemical vapor deposition

(PECVD) from trimethylsilane gas with tantalum nitride as adhesion promoter and

copper diffusion barrier [25], are also suggested for use as ILDs.

Other possible ILDs are porous silica xerogels, optically transparent

interconnected clusters of SiO2 of a small size, 20 nm, that may also serve as cladding

materials in optical interconnects, and suggested for use as alternatives to metal

interconnects. Copper adheres to α-SiOC:H films with thermal diffusion inhibition to ~

700 K [26]. SIOC films are promoted for use as intermetal dielectrics due to their

Material Dielectric constant (k)

Deposition method

Silicon dioxide (SiO2) 3.9-4.5

Fluorosilicate Glass (FSG) 3.2-4.0

Organosilicate Glass (SiOCH or SiCOH) 2.7-3.3

Fluorinated amorphous carbon (a-C:F) 2.4

PECVD

Parylene-N 2.7

Parylene-F 2.4-2.5 CVD

Polyimides 3.1-3.4

Fluorinated Polyimides 2.5-2.9

Benzocyclobutene polymers (BCB) 2.7

Hydrogen-Silsesquioxane (HSQ) 2.9-3.2

Methyl-Silsesquioxane (MSQ) 2.6-2.8

Poly(arylene ether) 2.8-2.9

Aromatic thermosets (SiLK) 2.6-2.7

Teflon AF 2.1

Porous silica (aerogel or xerogel) 1.1-2.4

Spin-on

Table 1.3 Dielectric constants of common dielectric materials [19]

(Used with permission, Semiconductor Fabtech).

11

higher thermal and mechanical stability as compared to organic films, and the ability to

tailor their density and refractive index by increasing their carbon content. However,

SIOC films take up oxygen radicals upon ashing [27]. Organosilicate glass (OSG) films

[28, 29] atomic layer deposition (ALD) Si:O:H films, mesoporous silica films, as well as

dual damascene structures of SiN/SiO2 have also been suggested.

VLK (very low k) films have been proposed as intra-layer dielectrics (trench level)

with SiO2 as the inter-layer (via level) dielectric, to reduce wiring capacitance and avoid

the lower thermal conductivity that results from full low-k (using low-k for both intra and

inter-layer dielectrics) [30]. In addition, HSQ, siloxane-based hydrogen silsesquioxanes

(HSiO3/2)2 have been proposed as an intermetal dielectric, with ammonia treatment to

form a nitride passivation layer inhibiting copper diffusion into these films [31].

1.5 Adhesion

In metallization, adhesion to the substrate is very important. The strength and

durability of an adhesive bond is significantly dependent on the interactions between the

adhering surfaces, functions of the chemical states, which can be monitored by XPS.

Good adherence means the film-substrate interface is not physically affected during

fabrication or expected use when exposed to reasonably low stress levels. Accepted

mechanisms of adhesive/adherend adhesion are based on mechanical interlocking,

surface energies, wetting adsorption, diffusion theory, and chemical bonding [32].

Wetting involves the establishment of intimate molecular contact described by

the contact angle, θ, of the liquid or molten species on the solid substrate, with smaller

12

contact angles corresponding to greater adherence, or wetting, between the film and the

substrate as a result of the interatomic-interactive forces acting across the film-substrate

interface. Therefore, intermolecular wetting is a prerequisite for developing strong

adhesion [32].

The mechanism of adhesion involves interactions of electrostatic forces or

chemical bonding. Atoms at the interface, held together by electrostatic forces when

ionized with opposite polarities, or polarization that creates dipoles, form attractions that

align to cause weak van der Waals type electrostatic bonding across the interface, that

are not strong enough to withstand dewetting. Larger polarizabilities and vibrational

frequencies increase the adhesion strength at the interface, resulting in partial wetting

Fig. 1.6 Wetting described by contact angle and temperature [34] (Used with permission, Kluwer Academic/Plenum Publishers).

Temperature

Con

tact

ang

le, θ

Dewetting

Partial Wetting

Complete Wetting0

13

[33]. Chemical bond formation at the interface is stronger than electrostatic or van der

Waals forces, resulting in much stronger adhesion across the interface and complete

wetting.

For example, Al wets SiO2 through bond formation from oxygen reduction,

whereas Cu may partially, or initially weakly, wet the surface due to van der Waals

forces. This is attributed to the relatively large difference in linear coefficient of thermal

expansion between Cu and SiO2 (Cu =16.5 x 10–6, SiO2 = 0.5 X 10–6). A strong bonding

force, typically supplied by an adhesion promoter, is needed to keep Cu films from

dewetting on SiO2.

Good adhesion requires strong inter-atomic bonding across the film-substrate

interface, low level of film stress, and absence of reactive environment. The more the

free energy is lowered, the stronger the adhesion. For adhesion agents to work they

must promote interaction between the interface surfaces. The mechanism of adhesion

involves the mechanical interlocking of the deposited metal with the polymer [35].

Titanium is often used as an adhesion promoter because it forms lower free

energy compounds with oxygen, nitrogen and carbon. Coupling agents, such as

organosilanes, are used to react with both the organic and inorganic surfaces by

forming a hydrophobic layer at the interface [36]. For polystyrene, this is done by

dissolving the coupling agent with the polymer molecules, e. g. silane, producing silane

anchoring sites that result in the polymerization of the molecule, with increased amounts

of silane producing increased adhesion [36]. Adhesion promotion is also accomplished

by irradiating individual polymer molecules with ultra-violet radiation [37], or electron

beams [38], or thermal curing [39], to polymerize the molecule by cross-linking.

14

In multilevel interconnect structures that use polymer dielectrics there are two

types of interfaces that have very different interface reactions, metal-on-metal and

polymer-on-metal. XPS studies have shown that reactive metals deposited on

polyimide (PI) break carbonyl bonds, however non-reactive metals such as Cu do not.

For copper, weak adhesion is often improved by plasma [40] or ion beam bombardment

[33], ammonia treatment of the polymer surface [32] or by inserting an adhesion layer

(Cr, Ti, Ta, TiN, TaN), depending on the properties of the metal/polymer interface at the

microscopic level.

Adhesion tests can be qualitative or quantitative. A simple qualitative test often

used to determine a minimum strength of adhesion between surface species is the

“Scotch tape test” [41-45]. In this test, a strip of tape is pressed on the top of the

deposited substance, and pulled back at approximately 180 ° in one smooth movement,

similar to a horizontal 180 º Peel Test shown in Figure 1.7 (P = applied load), and

examined for any film removed [46-48]. The European Metallizers Association

Technical Committee Films suggests the use of Scotch Tape 610 or equivalent with a

peel-force of approximately 3 N/15 mm measured against a steel-plate [49].

Fig. 1.7 Horizontal approximation of the 180 º Peel Test [46] (Used with permission, Elsevier).

P

P

15

Qualitative examination of the test substrate will reveal any metal transfer, with a strong

backlight facilitating visualization of any metal removed. The remaining exposed

surface is then analyzed for the deposited substance to see if any of it remains.

Quantitative measurements of adherence include the actual thickness of the deposited

film that remains after the “Scotch Tape Test” [46], and the peel-force when a calibrated

tape is used.

1.6 Techniques and Instrumentation

1.6.1 X-ray Photoelectron Spectroscopy

X-ray photoelectron spectroscopy (XPS) is a technique often used to analyze the

chemical properties of sample surfaces and thin films. In this technique x-rays are used

as a photon source to irradiate the sample to a depth of about 10 µm causing the

emission of electrons and any excess energy emitted as either fluorescence or, more

likely, as Auger electrons [50]. In XPS, photons with a minimum energy, hν, are

absorbed by atoms causing core electrons to be ejected from the sample. If the photon

energy, hν, is greater than the binding energy of the electron (EB), the excess energy is

the kinetic energy of the emitted photoelectron (EK). Calculation of the kinetic energy of

the emitted photoelectrons is the difference between the kinetic energy of the incident x-

rays, the measured binding energy and the work function, Φs, which is constant for that

specific spectrometer (Fig. 1.8)

EK = hν - EB - Φs

16

The binding energy of the photoelectron can be calculated from the measured

kinetic energy, K.E., work function, energy of the incident x-rays and any charging of the

sample surface, due to a sufficiently insulating sample that develops an imbalance of

charge and develops a potential difference relative to the spectrometer. This can be

calibrated by referencing peak energies relative to a calibrated internal standard e.g. C

1s 285.0 eV, to counteract charging effects [51, 52]:

EB = EK - hν- Φs - Ech

Typically, electrons with binding energies of 100 to 1000 eV are involved in these

transitions, with each chemical species producing electrons with unique energies, and

descriptive peak shapes, enabling XPS to differentiate between atoms with different

oxidation states e.g. Si 2+ and Si4+ by the shift in the peak binding energy called the

chemical shift. The chemical shift is based on the change in electrostatic potential on

the core electron when valence electron charge density is accepted or withdrawn from

Fig. 1.8 XPS electron emission process (Used with permission, Moulder & Froelich).

K (1s)

L1 (2s)

L2,3 (2p)

K (1s)

L1(2s)

L2,3 (2p)

hν

KE

KE

Initial state Final state Final stateK (1s)

L1 (2s)

L2,3 (2p)

K (1s)

L1 (2s)

L2,3 (2p)

K (1s)

L1(2s)

L2,3 (2p)

K (1s)

L1(2s)

L2,3 (2p)

hν

KE

KE

Initial state Final state Final state

17

the atom, enabling XPS for use in determining the chemical composition of surfaces and

the effects of processes on sample surfaces [50, 53]. Since binding energies of core

electrons are characteristic for elements in certain chemical environments, XPS also

enables determination of the electronic structure and band structure of interacting

species [50]. Observed chemical shifts facilitate descriptions of local coordination in a

system and the electronic change in oxidation state on adsorption, which is useful in

distinguishing adsorption sites of molecules adsorbed on surfaces, e.g. N2 (Fig. 1.9)

Fig 1.9 Nitrogen adsorbed on Ni(100) surface [54] (Used with

permission, Nilsson).

420 415 410 405 400 395

Binding energy (eV)

N 2 /Ni (100)

18

adsorbed on a Ni(100) surface. In Figure 1.9 the two N 1s peaks have a chemical shift

of 1.3 eV. The 399.4 eV peak corresponds to ionization of the outer N atom, and the

400.7 eV peak is due to ionization of the inner N atom [54].

XPS spectra can be used quantitatively because the intensity, I, of the peaks is

related to the atomic concentration of the species in the sample, N, the sensitivity of the

species to surface effects, A, as relative ratio [11, 50]:

ABB

BAA

AIN

AIN =

Changing the geometry of an instrument affects the relative intensities of different peaks

in analysis [55]; therefore, the intensity ratios should be calculated between results

obtained on calibrated instruments with the same geometric arrangement. XPS signals

measured at angles that are more surface sensitive (Figs. 1.10, 1.11), can be used to

Fig. 1.10 Comparison of angle resolved and normal emission geometry.

Mean Free PathMean Free Path

Surfacenormal

Escape Depthe - e-θ

Normal Emission, Angle Resolved,

19

approximate the relative thickness of a thin film, <100 nm, and indicate if a species is on

the surface or in the bulk of a sample, using the relationship between intensity and

thickness [38]:

)/()(

λdTiTi II −=

0

Auger spectra produced by XPS are the result of an x-ray induced photoelectron

ejection triggering an Auger electron emission from a valence energy level, and

consequent electron transfer from an outer energy level into the vacated position, core

hole, created by the photoelectron [50]. Auger electrons also give information about the

Fig 1.11 Angle resolved, θ = 15°, Si 2p spectrum enhances the surface oxide species as compared to Normal Si 2p spectrum, θ = 90° [56] (Used with permission, Briggs).

chemical species in the sample because the kinetic energy of the Auger electron is

directly related to the core binding energy of the photoelectron:

θ = 90°

θ = 15°

Binding Energy (eV)114 94

XPS

Inte

nsity

θ = 90°

θ = 15°

Binding Energy (eV)114 94

XPS

Inte

nsity

20

2,311 LLK2,3LKL EEEE −−=

Additional information about the type of species is given by the shape of the Auger

peak. Auger spectra are referenced to the vacuum level or the Fermi level, however the

standard vacuum level is defined as 4.500 eV above the Fermi level which is close to

the uncharacterized vacuum level that varies between instruments, so only the Fermi

level reference is absolute [57].

Figure 1.12 [58] schematically illustrates the basic parts of an x-ray photoelectron

spectroscopy instrument. The emitted electrons are routed through lenses, which

focus electrons upon the input slit, and retard electron kinetic energies, by a pre-set

retarding voltage, ER, collimated and focused on the entry slits of the cylindrical

Fig. 1.12 XPS Instrumentation [58] (Used with permission, Fairley).

x - ray gun

Samplestage

ion gun

slits

Host computer

Hemispherical Analyzer

Detector head

HT electronics

21

hemispherical analyzer. The hemispherical analyzer only allows electrons with

stipulated energies, determined by voltage parameters set for the lens system, to

proceed on to the detector. This energy, the pass energy, EP, is determined by the

operator and set in the parameters of the operational software.

The hemispherical analyzer and transfer lenses are generally operated in the

Fixed Analyzer Transmission (FAT) mode, also known as Constant Analyzer Energy

(CAE) mode. In CAE mode, electrons with kinetic energy, EK, lose energy by

retardation against an electrostatic potential, VR, prior to entry into the analyzer where

they pass with kinetic energy EP so that:

EK = e VR+ EP

Only electrons with EK = e VR+ EP will pass through the analyzer and be detected.

Thus, sweeping the electrostatic potential, VR, linearly with time allows the

measurement of the kinetic energy distribution of the photoelectrons emitted from the

sample:

EP = EK – ER

Whereas for Auger electrons the relationship is a ratio of electron energy and the

retardation ratio, R:

EP = EK /R

Most XPS spectra are acquired using CAE mode.

The resolution of XPS spectra is determined by the full width half maximum

(FWHM) of the detected lines affected significantly by the diameter of the analyzer, pass

energy and spread of energies in the x-ray source. In addition charging can interfere

22

with emitted energies causing broadening of the spectral lines. Non-monochromated x-

ray sources provide relatively large throughput to examine large surface areas [59].

1.6.2 Physical Vapor Deposition

Metallization of films is often accomplished by sputtering the surface of the film

with atoms of the desired metal or alloy. Physical vapor deposition (PVD) (Fig. 1.13)

has been widely used for this purpose by the semiconductor industry because most

films used in VLSI (very large scale interconnects) are alloys and PVD can provide

uniform alloy composition across the surface area of wafers while providing good step

coverage. Metal films commonly deposited by sputtering include Al-Cu, Ti-W, TiN, Pt,

and individual metals for specific purposes.

Fig. 1.13 Physical Vapor Deposition process [60] (Used with permission, University of Virginia).

23

For magnetron sputter deposition, an inert gas, often argon, is exposed to a glow

discharge created by applying dc voltage between the anode and cathode, at pressures

from 10-4 Torr to 1.0 Torr [61, 62]. The glow discharge ionizes the gas atoms and

accelerates them toward the metal target (cathode) with energies between 500 and

2000 eV [62], which then dislodge (sputter) metal atoms from the cathode. Sputtered

metal atoms with energies between 3 to 10 eV [62], leave the target, propelled toward

the sample attached to the anode, in a loop induced by a magnetic field. Metal atoms

with energies less than 5 eV [63] condense to form a thin film on the sample surface.

Those with energies > 5 eV can eject from the sample surface and readsorb on the

target surface. Contaminants in the PVD chamber are also excited by the impinging

argon ions, and atoms or ions with energies greater than 10 eV can damage the sample

surface creating defects, artifacts and induced extraneous reactions as shown in Figure

1.13 [60]. Therefore, the energy used is just enough to accomplish sputtering while

limiting extraneous interactions.

XPS is used to monitor the interaction of the deposited metal and the substrate

by plotting the ratio of the metal species intensity compared to the substrate intensity

against metal deposition time. Strong interactions between a deposited metal species

with free energy, γM, and the substrate, with free energy, γS, result in the metal

preferentially adhering to the substrate and conformal growth (Fig. 1.14(a)) for several

monolayers of deposition, termed Frank van der Merwe (F-M) growth, indicated by

several changes in slope of the metal deposition uptake curve.

γM < γS + γMS

24

Fig. 1.14 Three methods of metal film growth (a) Frank van de Merwe (b) Stranski - Krastanov (c) Volmer Weber

Stranski-Krastanov (S-K) growth begins with strong interactions between the deposited

metal species and the substrate that are overcome by the increase in interface energy

as the layer thickness increases due to strain to try to continue to fit the substrate. This

yields initial conformal growth followed by 3-D island formation indicated by a change in

the slope of the uptake curve (Fig. 1.14(b)). Weak interactions between the deposited

metal (M) and the substrate (S) result in the metal preferentially adhering to itself

resulting only in the formation of three dimensional islands (Fig. 1.14(c)), termed

Volmer-Weber (V-W) growth, indicated by a curve with an unchanging slope as

depicted in Figure 1.15.

(a)SPUTTER DEPOSITION

λMS > λM

Conformal growth

Free energy, λ, favors metal-substrate attraction

(c) λMS < λM

3-D islandingFree energy, λ, favors metal-metal attraction

Conformal growth with 3-D islanding

(b)

25

Fig. 1.15 Uptake curves of Metal/Substrate versus metal deposition time depicting two different growth modes: a) conformal growth b) 3-D island growth.

1.6.3 Precursor Deposition

Precursor deposition was accomplished by a simplified adaptation of Chemical

vapor deposition (CVD). CVD is a process involving the formation of a nonvolatile solid

film on a substrate by the reaction of vapor phase chemicals obtained from gases,

liquids, sublimable solids or combination materials that are stable at room temperature

and volatile with high partial pressure [33]. By definition the process introduces a given

composition of reactant gases and diluent inert gases into a reactant chamber; the gas

species move to the substrate and the reactants are absorbed on the substrate then the

adatoms undergo migration and film-forming chemical reactions after which the

gaseous by-products of the reaction are desorbed and removed from the reaction

M/S

XP

S In

tens

ity R

atio

M/S

XP

S In

tens

ity R

atio

0 2 4 6 8 10

Deposition Time or Coverage

No change in slope of curve

0 2 4 6 8 10

Deposition Time or Coverage

1st growth mode

2nd growth mode

0

2

4

6

8

0

2

4

6

8

(a) (b)

26

chamber [64]. In this simplified system the precursor is introduced into the reactant

chamber, without the need of a solvent. Deposition is carried out at a very low

temperature [65] to immobilize the species on the substrate surface and allow the

incorporation of species such as hydroxide or water [66], and enable UV or plasma

irradiation or ion bombardment, to inculcate the film with the desired properties. To get

good growth rates the reaction temperature should be lower than the melting point of

substrate and produce the desired species on the substrate with easily removable by-

products and low toxicity [67].

1.6.4. XPS, PVD and Precursor Doser Apparatus

The apparatus for physical and chemical vapor deposition combines the ultra-

high vacuum XPS system and DC magnetron PVD chamber depicted in Figure 1.16.

1.16 XPS, PVD and Precursor Doser System.

Residual Gas Analyzer

NH3 /H

2 O Doser

XPSPVD

Tower

SampleIntroduction

e- Beame- Beam

ion gunion gun

TurboPump

PVD Chamber(1 x 10-8 Torr)

Main Chamber(5 x 10-10 Torr)

Turbo Pump(backside)

VTMS Doser

Main Chamber(5 x 10-10 Torr)

27

The main chamber is equipped for CVD with a direct doser for film precursor

introduction, and a backfill doser for water and or ammonia introduction, ion-sputtering

gun for sample cleaning, a dual anode ray source and hemispherical analyzer for XPS

sample analysis. The PVD chamber consists of a water cooled sputter target about 10

cm from the sample and a magnetron power source monitored by a baratron gauge,

isolated from the main chamber by a gate valve.

28

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31

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33

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34

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35

CHAPTER 2

HYDROXYLATED AMORPHOUS POLYMERIZED CARBON:SILICON FILMS - ADHERENCE

/DIFFUSION BARRIERS FOR COPPER *

2.1 Motivation for Si:C:H Based Barrier Layer

The integration of Cu with SiO2 or lower dielectric constant materials in device

interconnect structures requires a barrier to inhibit Cu diffusion under bias/thermal

stress conditions due to copper’s relatively facile diffusion through SiO2 [1]. To date,

refractory metals such as Ta or Ta/TaN have been satisfactory in this regard, but the

move to porous low-k materials with relatively large coefficients of thermal expansion

has motivated a search for alternative barrier materials. Since many low-k materials

contain carbon, the formation of interfacial metal carbide layers, which may inhibit Cu

wetting [2], is an additional concern.

Recent work [3, 4] has shown that thin amorphous Si:C:H films (a-Si:C:H) exhibit

resistance to thermally stimulated Cu diffusion at temperatures below 800 K. More

recently, studies have shown that such films exhibit potentially suitable resistance to

combined bias/thermal stress (BTS) diffusion of copper [5], especially when partially

nitrided [6]. The use of thin a-Si:C:H films as surface modifications to intralayer low-k

dielectrics to act as Cu diffusion barriers, is therefore of significant interest [7].

*Entire chapter is reprinted from THIN SOLID FILMS, Vol 440, No 1-2, M. Pritchett et al., “Copper metallization of hydroxyl-modified amorphous Si:C:H films,” pp. 100-108, 2003 Elsevier Science B.V., with permission from Elsevier.

36

A major difficulty with the use of a-Si:C:H films in intralayer dielectric applications

is the lack of Cu/film adherence . When a-Si:C:H films are formed on various metal

substrates in the presence of electron or plasma bombardment, excellent adherence of

the film to the metal substrate is observed [3, 4, 8]. Subsequent deposition of Cu

adatoms onto the a-Si:C:H film, however, results in weak chemical interaction, and the

growth of three-dimensional islands (Volmer-Weber growth) [4, 9, 10], during the initial

nucleation process. Such weak Cu/film interaction results in facile thermal

agglomeration, and can result in void formation during via fill or device operation.

Wetting is the result of the attraction between the film and the substrate surface,

and is described thermodynamically as the free energy change of the system (∆γ) due

to the removal of a metal adatom from a metal cluster to the substrate surface:

∆γ = γF + γFS - γs (1)

Where γF, γS and γFS are the surface free energies of the metal film, substrate, and film-

substrate interface, respectively. When the overall change in free energy is negative,

the attraction of the film for the interface is greater than the attraction of the film particles

to each other, and wetting will occur. Since γF >>γs for many metal films on many

dielectric substrates, γFS must be large and negative for wetting to occur, i.e., there must

be a strong attractive chemical interaction between film and substrate.

If there is relatively little stress in the deposited film (e.g., due to lattice mismatch

with the substrate), layer-by-layer, or Frank-van de Merwe (FM) growth will result. More

commonly, lattice mismatch induces significant internal stress, and the increase in

stress with film thickness results in a transition from initial layer-by-layer growth to island

formation (Stranski-Krastanov, or SK growth) [11, 12]. Therefore, considerable insight

37

concerning the effects of substrate chemical structure on film nucleation and growth can

be gained by examining the interactions of the substrate with the first few monolayers of

deposited film [13].

Amorphous-Si:C:H films of controlled thickness and composition can be prepared

by electron bombardment of vinyltrimethylsilane (VTMS) multilayers under UHV

conditions [4]. Such films exhibit a chemical stoichiometry close to that of the VTMS

precursor [4] and display excellent thermal stability on a variety of substrates [3]. In

addition, VTMS films, polymerized by gamma radiation, have shown a decrease in their

gas permeability [14]. Other studies have demonstrated that additional properties of a-

Si:C:H films, such as leakage current and dielectric constant, can be adjusted by

modification of the films with ammonia [6], fluorine [15] and nitrogen [7]. The growth of

a-Si:C:H films can be adjusted by modification with oxygen [16]. Surface science

studies of the wetting of Cu on hydroxylated and dehydroxylated ordered alumina [17,

18] and hydroxylated Ta silicate [19] surfaces have found that hydroxylation promotes

the wetting of Cu. Therefore, modifying VTMS with H2O as a source of hydroxyl groups

might promote copper wetting on the modified a-Si:C:H film.

This work reports the ability of copper to wet, or grow conformally on, the surface

of a hydroxide modified cross-linked polymeric a-Si:C:H film (a-Si:C:H./OH) derived

from VTMS. At 90K physisorption of the VTMS precursor was immediately followed by

H2O physisorption. The addition of OH groups onto the Si:C:H film was achieved by

bombarding the physisorbed VTMS/H2O multilayers with an electron beam at 90K,

followed by annealing at 300K in UHV. XPS measurements indicate that Cu deposited

at 300 K by sputter deposition onto a-Si:C:H/OH grows conformally, with the initial Cu

38

ad-layer forming both Cu(I), from charge transfer between Cu and the Si:C:H film, and

Cu(0) species at the interface. In contrast, wetting and charge transfer were not

observed for Cu deposited onto a-Si:C:H, in agreement with previous results [3].

2.2 Experiment

The experiments were performed in an ultra-high vacuum (UHV) chamber that

has been described previously in detail [19]. Briefly, the system consists of a main

chamber with a base pressure of 6 x 10-10 Torr, and a dual-magnetron sputter

deposition chamber, with a base pressure of 5 x 10-9 Torr. Transferring the sample

between the two chambers was accomplished without atmospheric exposure. The

temperature of the sample was maintained between 90 K and 1200 K by a combination

of liquid nitrogen cooling and resistive heating.

X-ray photoelectron spectra were acquired in the UHV main chamber, using a

non-monochromated MgKα X-ray source at 15 kV and 300 W, with a commercial

hemispherical analyzer operated at 50 eV constant pass energy. Methods of data

acquisition and analysis have been described previously [19]. All spectra reported here

were acquired with the analyzer aligned along the surface normal (normal emission),

unless otherwise stated.

The Ti foil substrate had a purity of 99.97%, and was 0.0076 cm thick with an

area of 1 cm2. The foil was spot welded to Ti leads and a chromel-alumel thermocouple

was spot welded to the back of the sample. Sample cleaning was accomplished

through cycles of argon ion (Ar+) sputtering followed by annealing. The Ti substrate was

heated resistively through the thermocouple, and cooled indirectly by liquid nitrogen.

39

VTMS was procured from a commercial source, purified by freeze-thaw pumping, and

delivered through a stainless steel doser tube regulated by a manual leak valve, ~ 2 cm

from the face of the Ti sample. Gas exposures are reported here in Langmuir, L, (1L =

1 x 10-6 Torr-s), and have not been corrected for ion gauge sensitivity and flux to the

surface. To produce the unmodified film, 30L of VTMS was condensed onto a clean Ti

substrate at 90K, immediately followed by electron beam bombardment to induce

crosslinking [4, 20]. To produce the hydroxide-modified film, 30L of VTMS was

condensed onto a clean Ti substrate at 90K, immediately followed by 30L of H2O, and

electron beam bombardment.

A Kimball Electronics electron gun was used to produce an electron beam of 500

eV energy, with a 3mA emission current, producing an average flux at the sample

surface of 1.5 x 1013 electrons cm-2 s-1 [4]. After electron bombardment, the films were

annealed to 300K in UHV in order to desorb any remaining physisorbed species.

Copper was sequentially deposited by physical vapor deposition (PVD) onto the film’s

surface using an Advanced Energy MDX 500 magnetron source (25 watts), and an

MKS baratron gauge to monitor the Ar+ sputter pressure (2 x 10-5 atm), achieving a rate

of about 0.5 Å/minute, and sputtering Cu atoms with energies of ~ 0.2 eV to ~ 10 eV

[21, 22]. After copper deposition was completed, the sample was annealed in intervals

from 300K to 1000K by resistive heating. Copper and film interactions were

characterized by XPS after each successive deposition and temperature increment.

For both the unmodified and modified films, film compositions were calculated

using XPS core level data, with corresponding intensities, and appropriate elemental

40

sensitivity factors for this analyzer [23]. Relative film composition was determined from

XPS core level intensity data with the general equation:

NA/NB = [IAAB]/ [IBAA] (2)

NA and NB are the atomic concentrations of element A and B respectively, IA and IB are

the respective XPS intensities (proportional to peak areas), and AA and AB are the

appropriate atomic sensitivity factors [24].

The a-Si:C:H films produced were very thin, as demonstrated by the ability to

observe the titanium XPS peak after the 300K anneal, enabling the average thickness,

(d), of the a-Si:C:H films to be calculated as a function of the intensity of the Ti(2p3/2)

core level peak (ITi), and the inelastic mean free path (λ) of the Ti(2p3/2) photoelectron,

21.6 Å , using the following equation:

ITi = I Ti (0)exp(-d/λ) (3)

The thickness of copper deposited on the hydroxide-modified film was estimated from

the Si(2p) intensity using the intensity ratio of copper to silicon. The copper to silicon

intensity ratio, ICu/ISi, intensity ratio of reference element atomic sensitivity factors,

I∞Cu/I∞

Si, and the inelastic mean free paths for copper and silicon, (λCu) and (λSi), are

known, therefore the thickness, (dCu) is derived from the equation [24]:

ICu/ISi = (I∞Cu/I∞

Si)[(1-exp-dCu/λCu)/exp-dCu/λSi] (4)

Absolute copper and Ti intensities from clean samples in the system used for these

experiments have an estimated error in measurement of no more than 10% [4]. The

Scotch Tape Test was performed on the PVD copper film to assess adhesion [25, 26].

41

2.3 Results

2.3.1 Amorphous-Si:C:H Film Characterization

The average thickness of the unmodified a-Si:C:H cross-linked film produced

from the VTMS precursor, by the process described above, was determined (Eq. 3) to

be 15 Å. As shown in Figure 2.1, the Si(2p) XPS spectrum exhibits a peak at 100.8 eV,

-106 -104 -102 -100 -98 -960

1000

-290 -288 -286 -284 -282 -2800

4000 C(1s)

Si(2p)

XPS

Inte

nsity

(cps

)

Binding Energy (eV)

C - C, C - H

Si - C

Fig. 1 Fig. 1 Fig. 2.1 C(1s) and Si(2p) XPS Spectra for Unmodified a - Si:C:H film .

42

attributed to Si-C:H from VTMS [27], and a C(1s) spectra and a peak at 284.5 eV,

attributed to C-C and C-H bond types [27]. The unmodified a-Si:C:H film had negligible

oxygen content. No spectral changes were observed at glancing emission (not shown).

Although the use of extremely thin films (~ 30 Å) presents the risk of pinholes affecting

the results, Cu wetting results for unmodified a-Si:C:H (vide infra) were in agreement

with previously published results for thicker films [4].

Normal emission spectra for the hydroxide modified film, (a-Si:C:H/OH), are

shown in Figure 2.2. This film was determined to have an average thickness of 25 Å

(Eq. 3). The Si(2p) spectrum for the modified film displays a high binding energy

shoulder, while the C(1s) spectrum does not, indicating that the oxygen-containing

-544 -540-536-532-528-5240

1500

-290 -288 -286 -284 -282-2800

4500

-111-108-105-102-99 -96 -930

1000

C(1s)

O(1s)

Si(2p)

XPS

Inte

nsity

(cps

)

C -C, C -H

Si-O x

Si -OH

Si:C:H

Si-O xSi-OH

Fig. 2 Fig. 2

Binding Energy (eV)

Fig. 2.2 C(1s), Si(2p), O(1s) XPS spectra for the a-Si:C:H/OH film.

43

species are bonded primarily to silicon and not to carbon atoms. The O(1s) peak (Fig.

2.2) was fit with two spectral components with FWHM of 2.9 eV and binding energies of

532.3 eV for Si-Ox (x<2) and 533.7 eV for Si-OH [28]. The Si(2p) spectrum (Fig. 2.2),

exhibits a dominant feature at 100.8 eV binding energy for Si-C:H. The other features

are more difficult to assign due to the variety of assignments found in the literature for

Si-OH. The shoulder at higher binding energies is well fit by two peaks, with a

component at 102.3 eV peak assigned to Si-Ox (x<2) as indicated by Hollinger [29] and

Himpsel [30], and a component at 103.8 eV attributed to a hydroxylated silicon oxide

species [31]. A binding energy of 102.3 eV, however is also consistent with a

hydroxylated Si+1 species, Si-OH [32, 33]. Evidence for a hydroxylated Si species is

more clearly indicated by the 533.7 eV peak in the O(1s) spectrum, which is in

agreement with the assignment of 533.5 eV for the O(1s) Si-OH peak [28]. The

absence of significant numbers of C-OH groups is indicated by the absence of a

characteristic peak near 286.0 eV in the C(1s) spectrum [27].

Calculation of the C(1s)/Si(2p)total atomic ratio, using appropriate sensitivity

factors (A), and inelastic mean free path (λ) values, atomic concentrations(N), and peak

intensities(I) [23]: yielded a C(1s)/Si(2p)total ratio of 3.5:1, for a-Si:C:H/OH, which is in

agreement with previous research that calculated the C(1s)/Si(2p)total ratio for a-Si:C:H

as 3.8:1 [4]. This prior work found that chemical and matrix effects, due to differences

in kinetic energies, could be corrected for by dividing atomic concentrations by their

appropriate mean free paths as reported by Ke, et al. [24, 34]:

NC / NSi = [(IC · ASi) / λC ] / [(ISi · AC) / λSi] (5)

44

This treatment was applied to these results using inelastic mean free paths calculated

from the modified form of the Bethe equation [35]: C(1s) = 22.5 Å and Si(2p) = 25.7 Å

and O(1s) = 15.0 Å. The corrected C(1s)/Si(2p)total ratio for both the unmodified a-

Si:C:H film and the modified a-Si:C:H/OH film increased to ~ 4.5:1. Because of the

uncertainty of the actual additional chemical and matrix effects contributed by oxygen in

a-Si:C:H/OH, a comparison calculation was done using only the a-Si:C:H Si(2p)(101)

peak intensity to calculate the C(1s)/Si(2p)total ratio, which resulted in a C(1s)/Si(2p)total

ratio of 4.7:1. Both methods of calculating the C(1s)/Si(2p)total ratio for a-Si:C:H/OH,

correlate well with the five to one C(1s)/Si(2p)total stoichiometry of VTMS. The

O(1s)total/Si(2p)total atomic ratio for the same a-Si:C:H/OH film was calculated to be

0.44:1, and 0.073:1 for the unmodified a-Si:C:H film.

2.3.2 Copper Deposition by Physical Vapor Deposition (PVD)

XPS was used to characterize the interactions of sputter-deposited (often

referred to as physical vapor deposited--PVD) Cu with clean and modified film

substrates. Cu coverage was ascertained from Cu(2p3/2)core level spectra, with Cu(I)

and Cu(0) oxidation states distinguished by the examination of x-ray excited Cu(L3VV)

Auger spectra [17]. Cu(2p3/2) spectra at no time revealed observable satellite peaks,

indicating the absence of significant concentrations of Cu(II) [36, 37]. Copper was

deposited in the sputter deposition chamber at 300K at intervals of several minutes

sputter deposition time, with estimated kinetic energies between ~ 0.2 eV and ~ 10 eV

[21, 22]. XPS data were acquired after each deposition, followed by additional copper

deposition. Energetic particles, such as those that can be deposited during PVD, are

45

known to deform and modify surfaces [38]. In this case, however particle energies do

not seem to be the cause of the chemical modification of the modified a-Si:C:H film that

enabled wetting, since copper PVD on unmodified a-Si:C:H films using the same

parameters, did not produce similar XPS spectra or copper wetting.

The evolution of the Si(2p) spectrum during Cu deposition is shown in Figure 2.3

for the unmodified a-Si:C:H film, and in Figure 2.4 for a-Si:C:H/OH. No changes were

observed for C(1s) spectra (not shown) after deposition on either modified or

in binding energy or intensity with Cu deposition (Fig. 2.3). In contrast, Cu deposition

Fig. 2.3 Fig. 2.4

-105 -103 -101 -99 -97

-105 -103 -101 -99 -97

-105 -103 -101 -99 -97

-105 -103 -101 -99 -97 Binding Energy (eV)

XPS

Inte

nsity

( cp

s)

Cop

per D

epos

ition

tim

e (m

in)

Si(2p)

Fig. 3Fig. 3

- 0.0

- 1.0

- 3.0

- 6.0

Fig. 2.3 Si(2p) XPS spectra during deposition of a-Si:C:H film.

-111-108-105-102-99 -96 -93

-111-108-105-102-99 -96 -93

-111-108-105-102-99 -96 -93

-111-108-105-102-99 -96 -93

- 6.5

- 2.5

- 8.5

- 0.0

Binding Energy (eV)

XPS

Inte

nsity

( cp

s)

Si(2p)

Fig. 4Fig. 4

Cop

per D

epos

ition

tim

e (m

in)

Si:C:H

Si:C:H

Si:C:H

Si:C:H

Fig. 2.4 Si(2p) XPS spectra during

Si:C:H

Si:C:H

Si:C:H

Si:C:H

Si-OH

Si-OH

Si-OH

Si-OH

deposition of a-Si:C:H/OH film.

46

on the hydroxylated substrate results in relative attenuation of the Si(2p) intensity

attributed to Si-C:H sites (Fig. 2.4). Changes in relative Cu(2p3/2)/Si(2p) intensity with

Cu deposition time are plotted in Figure 2.5. The plot of Cu(2p3/2)/Si(2p) intensity vs. Cu

deposition for unmodified VTMS exhibits linear behavior, which is consistent with 3-

dimensional islanding (Volmer-Weber growth) [11]. This contrasts with the behavior

observed on OH-modified substrates-- a linear increase followed by a distinct change in

slope, which is more similar to Stranski-Krastanov growth [11].

The evolution of the x-ray excited Cu(L3VV) spectrum during deposition on

unmodified and on hydroxylated films is shown in Figure 2.6. For the first six minutes of

copper deposition on the hydroxylated film at 300K (Fig. 2.6a), the Cu(L3VV) Auger

0 2 4 6 8 10 12 0 2 4 6 8

10 12 14 16 18 20

(a)

(b)

Cu(0)

Cu(I)

Break Point

Deposition Time (min)

XPS

Inte

nsity

Rat

io C

u(2p

3/

2 )/ (S

i 2p)

Fig. 2.5 Copper uptake curve for copper deposition on a) unmodified a-Si:C:H film b) OH-modified a-SI:C:H film.

47

spectrum remains broad, with features at 914.9 eV and 918.0 eV, consistent with the

formation of both Cu(I) and Cu(0) [39, 40]. (Cu(II) was eliminated as a possible source

of the 918.0 eV peak because there were no characteristic satellite peaks in the

Cu(2p3/2) spectra--not shown). At longer deposition times, the Cu(L3VV) peak

noticeably narrows and shifts to 918.8 eV, indicating the predominance of copper metal

[39, 40].

Fig. 2.6 Cu Auger spectra for copper deposition on a) OH-modified a-Si:C:H film b) unmodified a-Si:C:H film.

Deposition on the unmodified a-Si:C:H film (Fig. 2.6b) results in the exclusive

formation of Cu(0). The feature in Figure 6b, assigned to Cu(0), is observed near 916

eV for short Cu deposition times, and is shifted to higher binding energies upon

additional deposition. Such a shift is characteristic of the growth of Cu(0) clusters [41]

and is due to the enhanced screening of Auger final state holes with increasing cluster

890 900 910 920 930

a-Si:C:H Film

Aug

er In

t ens

ityDeposition Time (min)

Kinetic Energy (eV)900 910 920 930

0

500

1000

1500

2000

2500

3000

3500

a-Si:C:H/OH Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

0.0

7.5

3.5

2.0

1.0

5.5

8.5

Cu(0)Cu(I) Cu(0)

20.5

14.5

8.5

6.53.51.0

890 900 910 920 930

a-Si:C:H Film

Aug

er In

t ens

ityDeposition Time (min)

Kinetic Energy (eV)900 910 920 930

0

500

1000

1500

2000

2500

3000

3500

a-Si:C:H/OH Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

0.0

7.5

3.5

2.0

1.0

5.5

8.5

Cu(0)Cu(I) Cu(0)

20.5

14.5

8.5

6.53.51.0

48

size [42]. There is also a shift to lower binding energy in the Cu(2p3/2) spectrum (not

shown) that has been reported for the formation of Cu clusters [43].

For Cu deposition on the modified film, the narrowing of the Cu(L3VV) Auger

peak shape in increasing Cu coverage, accompanied by the shift to higher kinetic

energy (Fig. 2.6a) marks the evolution of a dominant Cu(0) feature that correlates with

the break point in the Cu(2p3/2)/Si(2p) uptake curve (Fig. 2.5b), and is indicative of

copper wetting, or SK growth, on the modified a-Si:C:H/OH film. This behavior is not

observed for Cu deposition on the unmodified a-Si:C:H film, which produces no change

in the uptake slope, consistent with Volmer-Weber growth (3-D islanding) [9], and Cu(0)

cluster formation [39, 40]. An examination of Figure 6a, however, reveals that Cu

deposited at coverages before the break point, (Fig. 2.5b), includes both Cu(I) and

Cu(0). At higher coverages, only Cu(0) is formed. After the Scotch Tape Test, XPS

characterization reveals the presence of Cu.

2.3.3 Annealing Effects

In order to assess the stability of the Cu/substrate interface for the hydroxylated

film, Cu was deposited on a hydroxylated substrate at 300 K, followed by 15 minute

anneals at elevated temperatures in UHV, with XPS acquisition during cool down.

Results are shown in Figure 2.7 for a copper film with an initial average thickness of 5 Å

prior to annealing. The Cu(2p3/2)/Si(2p) ratio decreases slightly due to island formation

between 500K and 800K, after which there is a much more significant decrease in

relative Cu intensity. Such a decrease can be caused either by island formation

(dewetting) or by diffusion of Cu from the surface into the bulk. A measurement of the

49

relative intensities of the Cu(2p3/2)and Cu(3p) intensities can distinguish between these

two possibilities. Since the inelastic electron mean free paths of the two copper species

in the film are very different, 13.9 Å for Cu(2p3/2) and 38.2 Å for Cu(3p) , the Cu(2p3/2)

peak intensity will be preferentially attenuated relative to the Cu(3p) intensity if copper

diffuses into the bulk. Due to the limitations of XPS sensitivity, such a method is

obviously less sensitive than electrical measurements for the detection of Cu diffusion,

but does serve to delineate the reason for changes in relative Cu XPS intensity. A plot

of the Cu(3p)total/Cu(2p3/2)total intensity ratio against temperature can therefore

distinguish between copper island formation and copper diffusion, because copper

200 400 600 800 1000

0

1

2

3

4

5

0.0

0.2

0.4

0.6

0.8

1.0

1.2

Temperature (K)

(b) X

PS In

tens

ity C

u32p

)/Cu(

2p),

cps

(a)

XPS

Inte

nsity

Cu(

2p)/S

i(2p)

, cps

Fig. 7 Fig. 7

(a) (b)

diffusion

Fig. 2.7 Dewetting is indicated by the decrease in curve (a) after 500 K.

Diffusion is indicated by the increase in curve (b) after 800 K.

Diffusion indicated

50

diffusion will produce a significant decrease in the Cu(2p3/2) peak intensity and a

subsequent increase in the Cu(3p)/Cu(2p3/2) intensity ratio [44]. The data in Figure 2.7

indicate no significant change in the Cu(3p)/Cu(2p3/2) intensity ratio below 800 K.

Above 800 K the ratio increases sharply, indicating that annealing to temperatures of

800 K or higher induces considerable Cu diffusion into the substrate. Similar results

were observed for Cu deposited on unmodified a-Si:C:H [3]. Annealing above 800 K

also produced significant changes in the silicon XPS spectrum of the hydroxide modified

a-Si:C:H film (not shown). The dominant Si(2p) 100.8 eV peak (Si-C:H) at 300K, began

to significantly decrease in intensity after 800K, until only 10% of its intensity remained

at 1000K. Annealing above 800K also produced two new Si(2p) peaks. The 98.1 eV

peak is attributable to Ti silicide formation, and the 99.6 eV peak indicates Si-Si

bonding.

2.4 Discussion

The data presented here demonstrate that hydroxylation of the surface of an a-

Si:C:H film strengthens the Cu/film chemical interactions and results in the deposition of

a conformal first layer of PVD Cu at 300 K. The evolution of a dominant Cu(0) feature

for deposition on the hydroxylated film (Fig. 2.6a) correlates with the break point in the

Cu(2p3/2)/Si(2p) uptake curve, depicted in Figure 2.5b. Such a break point is indicative

of initial conformal growth on the modified a-Si:C:H/OH film [11]. Unlike many cases of

S-K growth [17], however, the first layer of deposited Cu consists of both Cu(I) and

Cu(0) species.

51

Different behavior is observed for Cu deposition on an unmodified a-Si:C:H film,

where no change in the uptake slope is observed (Fig. 2.5), consistent with Volmer-

Weber growth (3-D islanding) [9, 11], and for which only Cu(0) formation is observed

(Fig. 2.6b). The shift of the Cu(L3VV) peak to higher Auger kinetic energy with

increasing Cu deposition time (Fig. 2.6b) is consistent with the growth in size of Cu(0)

clusters, since increasing cluster size results in more effective screening of Auger final

state holes [39, 40]. The conclusion of conformal Cu growth on the OH-modified film

surface is further supported by the fact that the substrate Si(2p) intensity is significantly

attenuated during copper deposition on the OH-modified film, which is expected for

conformal growth (Fig. 4) due to chemical interactions of the Cu adatoms with the

modified substrate. Three dimensional island growth of Cu, on the other hand, will not

significantly attenuate the substrate Si(2p) signal as shown in Fig. 3.

Additional information about the copper-film interface can be obtained from the

Auger parameter, α, defined by the equation:

α = EKCu(L3VV) + EBCu(2p3/2) (7)

since the value of α is independent of the uniform sample charging commonly

encountered for XPS on insulating samples [43]. At very low Cu coverage (< 1 minute

deposition time, ~ 0.25 ML) on a-Si:C:H/OH, a broad x-ray excited Cu(L3VV) feature is

observed (Fig. 2.6a). The two partially resolved peaks have α values of 1848.2 eV and

1850.6 eV (Fig. 2.8). The former value is consistent with Cu(I) formation [39, 45],

although the latter is low for the formation of Cu(0) species, which may be due to cluster

formation [40, 45, 46]. At two minutes deposition time (a Cu coverage of ~ 0.5 ML) the

52

first Auger parameter (corresponding to Cu(I)) does not change significantly, but the

second Auger parameter shifts to 1851.5 eV, which is now in the range generally

reported for Cu(0) [45, 47].

Fig. 2.8 Cu Auger parameter indicates both Cu(I) and Cu(0) species are deposited.

In contrast, Cu deposition on a-Si:C:H results in a similar limiting value for α only

after much longer deposition times (> 8.5 minutes). Comparisons with core-hole

spectra (not shown) indicate that changes in α are due mainly to changes in the Cu(L3

VV) spectra (Fig. 2.6a, 6b), and therefore to changes in final state screening [39, 40,

43]. The fact that an α value characteristic of bulk Cu is observed after only 2.5 minutes

Cu deposition on the hydroxylated surface, but not until after 8.5 minutes on the

unmodified surfaces, indicates a more rapid coalescence of isolated Cu clusters on the

1847.1848.1848.1849.1849.1850.1850.1851.1851.1852.1852.1853.

0 1 2 3 4 5 6 7 8 9

Cu Deposition time (min)

Aug

er P

aram

eter

(α

')

Cu +1

Cu 0

Cu Deposition Time (min)

Aug

er P

aram

eter

(α)

53

hydroxylated surface. This is also consistent with the more rapid attenuation of the

substrate Si(2p) intensity on the hydroxylated surface, and confirms the tendency

toward conformal Cu growth on a-Si:C:H/OH.

The thickness of copper deposited on a-Si:C:H/OH was estimated from the

Si(2p) intensity using the intensity ratio of copper to silicon film substrate (Eq. 4). The

average thickness at the break point in the Cu/Si uptake curve (Fig. 2.5) is 2 Å, which is

very close to the Cu(I) ionic diameter of 1.92 Å [40]. Since the initial deposition of Cu is

marked by the formation of both Cu(I) and Cu(0), it is difficult to give a precise

determination of surface coverage in the first layer, the more so as the film surface is

undoubtedly far from smooth on an atomic scale, but these data indicate that the

coverage of the initial Cu adlayer is slightly less than 1 monolayer. The data in Figures

2.4 - 2.6 therefore indicate that Cu deposition at 300 K on the OH-modified a-Si:C:H film

results in initial conformal growth, with the first layer including both Cu(I) and Cu(0)

species. Upon completion of the first layer, a second, metallic Cu layer begins to form.

In contrast, Cu deposition at 300 K on the unmodified a-Si:C:H film results exclusively in

Cu(0) formation and VW growth (non-wetting) as shown in Figures 5a and 6b. Similar

hydroxyl-induced transitions from non-wetting to wetting behavior for deposited Cu have

been reported for sapphire(0001) surfaces [17], and for Rh deposited on ordered,

ultrathin alumina films [48]. Cu wetting has been reported for deposition at 300 K on

hydroxylated ultrathin Ta silicate films on Si(100) and for air-exposed TaSiN substrates

[49]. Theoretical studies also indicate that Pt nucleation occurs preferentially at surface

OH sites on MgO single crystal surfaces [50]. These results strongly suggest, therefore,

54

that surface hydroxylation may be an effective method for control of the nucleation of

semi-noble metal adatoms on a broad variety of oxide or other dielectric substrates.

Thermally-induced copper diffusion was observed by XPS on OH-modified a-

Si:C:H films only at temperature above 800K in UHV, as shown by the lack of change in

the Cu(3p)/Cu(2p3/2) versus temperature curve (Fig. 2.7). The thermal stability of the

Cu/substrate interface is similar for OH-modified and unmodified a-Si:C:H films [4].

Annealing above 800K caused copper to diffuse into the film and resulted in the

development of larger clusters of Cu(0) [43], at the expense of Cu(I), as indicated by the

peak shift from 918.0 eV to higher kinetic energy, 918.6 eV, corresponding to an Auger

parameter shift in the range for Cu(0), from 1851.3 eV to 1851.9 eV.

In addition, annealing from 500K to 800K resulted in a decrease in Si(100.8 eV)

peak intensity, and increases in Si(102.3 eV) and Si(103.8 eV) peak intensities, which

may be due to uncharacterized reactions within the a-Si:C:H/OH film. Annealing to

1000K resulted in a large decrease in intensity of the Si 100.8 eV, Si 102.3 eV and Si

103.8 eV peaks, and the production of a large peak at 98.1 eV, likely titanium silicide,

and a small 99.6 eV peak, attributed to Si-Si, indicating decomposition of the film.

2.5 Conclusions

This study demonstrates that hydroxyl modification of vinyltrimethylsilane (VTMS)

derived amorphous silicon:carbon:hydrogen (a-Si:C:H) film facilitates the wetting of

sputter deposited, physical vapor deposited (PVD), copper to the hydroxyl modified a-

Si:C:H/OH film’s surface at 300K. This study also shows that modification of a-Si:C:H

55

with hydroxide, does not adversely affect the film’s ability to inhibit copper diffusion to

800K, as characterized by x-ray photoelectron spectroscopy (XPS).

These findings suggest the use of a-Si:C:H/OH films as a combined adhesion

promoter and diffusion barrier for the deposition of the initial copper seed layer in

interconnect circuit development, and indicate the possibility that additional

modifications of these films, to lower their dielectric constant and enhance other

properties, may be possible without altering the adhesion/barrier properties, enabling

their use as intermetal dielectrics that do not require an additional diffusion barrier layer

or adhesion promoter.

56

2.6 Chapter References

[1] McBrayer, J. D., Swanson, R. M. and Sigmon, Y. W., J. Electrochem. Soc. 133

(1986) 1242-6.

[2] Chen, L., Magtoto, N., Ekstrom, B. and Kelber, J., Thin Solid Films 376 (2000)

115-123.

[3] Chen, L. and Kelber, J. A., J. Vac. Sci. Technol. A 17 (1999) 1968.

[4] Tong, J., Martini, D., Magtoto, N., Pritchett, M. and Kelber, J., Appl. Surf. Sci. 187

(2002) 256-264.

[5] Lanckmans, F., Gray, W. D., Brijs, B. and Maex, K., Microelectron. Eng. 55

(2001) 329-335.

[6] Ishimaru, T., Shioya, Y., Ikakura, H., Nozawa, M., Nishimoto, Y., Ohgawara, S.

and Maeda, K., Proceedings of the International Interconnect Technology

Conference, San Francisco, CA, 2001, Institute of Electrical and Electronics

Engineers, Inc., (2001) 36-38.

[7] Lee, S. G., Kin, Y. J., Lee, S. P., Oh, H.-S., Lee, S. J., Kim, M., Kim, I.-G., Kim,

J.-H., Shin, H.-J., Hong, J.-G., Lee, H.-D. and Kang, H.-K., Jpn. J. Appl. Phys. 40

(2001) 2663-2668.

[8] Deplancke, M. P., Powers, J. M., Vandentop, G. J., Salmeron, M. and Somorjai,

G. A., J. Vac. Sci. Technol. A 9 (1991) 450.

[9] Aubin, E. and Lewis, L. J., Phys. Rev. B 47 (1993) 6780-6783.

[10] Miller, T., Rosenwinkel, E. and Chiang, T.-C., Phys. Rev. B 30 (1984) 570-577.

[11] Rhead, G. E., Barthès, M.-G. and Argile, C., Thin Solid Films 82 (1981) 201.

57

[12] Komori, F., Kushida, K., Hattori, K., Arai, S. and Iimori, T., Surf. Sci. 438 (1999)

123.

[13] Varma, S. and Chottiner, G. S., J. Vac. Sci. Technol. 10 (1992) 2857-2862.

[14] Yampol'skii, Y. P. and Volkov, V. V., J. Membr. Sci. 64 (1991) 191-228.

[15] Yoshimaru, M., Koizumi, S. and Shimokawa, K., J. Vac. Sci. Technol. A 15

(1997) 2915-2922.

[16] Lee, M.-S. and Bent, S. F., J. Vac. Sci. Technol. A 16 (1998) 1658-1663.

[17] Niu, C., Shepherd, K., Martini, D., Tong, J., Kelber, J. A., Jennison, D. R. and

Bogicevic, A., Surf. Sci. 465 (2000) 163-176.

[18] Heemeier, M., Frank, M., Libuda, J., Wolter, K., Kuhlenbeck, H., Baumer, M. and

Freund, H. J., Catalysis Letters 68 (2000) 19-24.

[19] Shepherd, K. and Kelber, J., Appl. Surf. Sci. 151 (1999) 287-298.

[20] Toth, A., Bertoti, I., Marletta, G., Ferenczy, G. and Mohai, M., Nucl. Instrum.

Methods Phys. Res., Sect. B 116 (1996) 299-304.

[21] Wasa, K. and Hayakawa, S., in R. F. Bunshah and G. E. McGuire (Eds.):

Handbook of Sputter Deposition Technology, Noyes Publications, Park Ridge,

New Jersey 1992, p. 32.

[22] Zou, W., Ph. D. Dissertation, University of Virginia, Charlottesville, Virginia 2001,

p. 303.

[23] "Vacuum Generators", Hastings, UK,

[24] Seah, M. P., in D. Briggs and M. P. Seah (Eds.): Practical Surface Analysis, Vol.

1, New York 1990, p. 201-255.

58

[25] Chang, C.-A., Baglin, J. E. E., Schrott, A. G. and Lin, K. C., Appl. Phys. Lett. 51

(1987) 103.

[26] Chan, C. M., Cao, G. Z., Fong, H., Sarikaya, M., Robinson, T. and Nelson, L., J.

Mater. Res. 15 (2000) 148-154.

[27] Wagner, C. D., Naumkin, A. V., Kraut-Vass, A., Allison, J. W., Powell, C. J. and

Jr., J. R. R., www.nist.gov, "NIST X-ray Photoelectron Spectroscopy Database",

NIST, 2002.

[28] Miller, M. L. and Linton, R. W., Anal. Chem. 57 (1985) 2314-19.

[29] Hollinger, G., Appl. Surf. Sci. 8 (1981) 318-36.

[30] Himpsel, F. J., McFeely, F. R., Taleb-Ibrahimi, A. and Yarmoff, J. A., Phys. Rev.

B 38 (1988) 6084-6096.

[31] Barr, T. L., J. Phys. Chem. 82 (1978) 1801-1810.

[32] Fujita, K. and Oya, A., J. Mater. Sci. (UK) 31 (1996) 4609-4615.

[33] Iiojima, Y. and Sonoda, N., Clean Technol. 12 (1992) 39.

[34] Ke, R., Haasch, R. T., Finnegan, N., Dottl, L. E., Alkire, R. C. and Farrell, H. H.,

J. Vac. Sci. Technol. A 14 (1995) 80-88.

[35] Powell, C. J., Jablonski, A., Tilinin, I. S., Tanuma, S. and Penn, D. R., J. Electron

Spectrosc. Relat. Phenom. 98-99 (1999) 1-15.

[36] Chusuei, C. C., Brookshier, M. A. and Goodman, D. W., Langmuir 15 (1999)

2806-2808.

[37] Paszti, Z., Peto, G., Horvath, A. E. and Karacs, A., J. Phys. Chem. B 101 (1997)

2109-2115.

[38] Paterson, M. J., Diamond Relat. Mater. 5 (1996) 1407-1413.

59

[39] Klein, J. C., Proctor, A., Hercules, D. M. and Black, J. F., Anal. Chem. 55 (1983)

2055-2059.

[40] Kowalczky, S. P., Pollak, R. A., McFeely, F. R., Ley, L. and Shirley, D. A., Phys.

Rev. B 8 (1973) 2387-91.

[41] Wertheim, G. K. and DiCenzo, S. B., Phys. Rev. B 37 (1988) 844-847.

[42] Jennison, D. R. and Madden, H. H., Phys. Rev. B 21 (1980) 430-435.

[43] Wu, Y., Garfunkel, E. and Madey, T. E., J. Vac. Sci. Technol: A 14 (1996) 1662-

1667.

[44] Chen, J. G., Colaianni, M. L., Weinberg, W. H. and J.T. Yates, J., Surf. Sci. 279

(1992) 223-232.

[45] Moretti, G., J. Electron Spectrosc. Relat. Phenom. 95 (1998) 95-144.

[46] Mason, M. G., Phys. Rev. B 27 (1983) 748-762.

[47] Moulder, J. F., Stickle, W. F., Sobol, P. E. and Bomben, K. D., Handbook of X-

ray Photoelectron Spectroscopy, Physical Electronics, Eden Prairie, MN, 1995.

[48] Libuda, J., Frank, M., Sandell, A., Andersson, S., Bruhwiler, P. A., Baumer, M.,

Martenson, N. and Freund, H.-J., Surf. Sci. 384 (1997) 106.

[49] Zhao, X., Leavy, M., Magtoto, N. P. and Kelber, J. A., Appl. Phys. Lett. 79 (2001)

3479-3481.

[50] Bogicevic, A. and Jennison, D. R., Surf. Sci. Lett. 437 (1999) L741-L747.

60

CHAPTER 3

NITROGEN MODIFIED HYDROXYLATED AMORPHOUS POLYMERIZED

CARBON:SILICON FILMS - ADHERENCE /DIFFUSION BARRIERS FOR COPPER

3.1 Motivation for Nitrogen Modification of a-Si:C:H Films

More densely integrated circuits require a change from aluminum to copper

metallization and a concomitant change to lower dielectric constant (low-k) interconnect

materials. Since copper diffuses into SiO2 and low-k hydrogen silsesquioxane (HSQ)

films [1] under bias/thermal stress (BTS) conditions [2], low-k barrier layers, or

dielectrics, must be developed that will inhibit diffusion into SiO2 [3]. In addition copper

must adhere to, or wet, these barrier layers.

Results detailed in Chapter 2 show that hydroxylated amorphous silicon carbon

hydrogen films (a-Si:C:H/OH) not only resist thermally stimulated copper diffusion at

temperatures up to 800K [4, 5], but also enable copper wetting of these hydroxylated

films [6-8]. Additional modifications of these films with nitrogen may produce films that

possess low dielectric constants [9], while acting as barrier layers to copper diffusion

and promoting copper wetting, suggesting their use as all in one intermetal dielectric

layers that are potentially resistant to combined bias/thermal stress diffusion of copper

[10]. Therefore, the modification of amorphous silicon carbon hydrogen (a-Si:C:H) films

61

with nitrogen, as well as the modification of a-Si:C:H films with hydroxide, is of interest

in determining their use as intermetal low-k dielectrics [9].

The first requirement for nitrogen modified a-Si:C:H films as dielectrics, is that

copper wet, grow conformally, on the surface of these films. The second requirement is

that these nitrogen modified a-Si:C:H films inhibit copper diffusion while adhering to the

substrate surface. It has been previously determined that unmodified and hydroxide

modified a-Si:C:H films, formed on a titanium substrate in the presence of electron or

plasma bombardment, have excellent adhesion of the film to the metal substrate [4, 5,

8, 11]. Subsequent deposition of Cu adatoms onto the surface of hydroxyl modified a-

Si:C:H films results in conformal growth, wetting [8], in contrast with 3-D islanding

observed for unmodified a-Si:C:H films [5, 8, 12]. The question now is how modifying a-

Si:C:H films with nitrogen will affect the adhesion and barrier properties of these films.

Nitrogen is an electronegative element that may affect the attractive forces at the

interface of the film and the substrate, previously described thermodynamically as the

change in free energy change of the system (∆γ). Since wetting occurs when film

particles preferentially attract the copper adatoms, it may be that nitrogen will act to

enhance the ability of copper to lose electrons as compared to the unmodified film. This

may resulting in an increased attraction between the film particles and the copper

adatoms [13] such that the free energy of the film-surface interface, γFS, mitigates the

free energy of attraction between film particles, γF, and in conjunction with the free

energy of the substrate, γS, produces a negative overall change in free energy:

∆γ = (γF + γFS) - γS (1)

62

A strong attraction between the nitrogen modified film and copper adatoms, as was

found for the hydroxylated a-Si:C:H film and copper, would indicate that copper wetting

of the film’s surface is unaltered by the incorporation of nitrogen into the a-Si:C:H film.

Nitrogen modified amorphous Si:C:H films with a chemical stoichiometry close to

the C5H12Si of the precursor Vinyltrimethylsilane (VTMS), with controlled composition

and thickness, were prepared under UHV conditions [8],in anticipation of positively

altering film properties such as leakage current [14], and dielectric constant [9, 15]. This

chapter reports the ability of copper to wet, or grow conformally on, the surface of

nitrogen and nitrogen and hydroxyl modified cross-linked polymeric a-Si:C:H films

derived from VTMS, as compared to unmodified and hydroxyl modified a-Si:C:H films.

At 90K, in UHV, the VTMS precursor was physisorbed on titanium, immediately

followed by electron bombardment and annealing to 300K in UHV, resulting in an

unmodified amorphous Si:C:H film (a-Si:C:H). The nitrogen modified film was formed at

90K, in UHV, by condensing the VTMS precursor on the titanium substrate, immediately

followed by ammonia physisorption, electron bombardment, and annealing to 300K,

resulting in a nitrogen modified amorphous Si:C:H film (a-Si:C:H/N). The nitrogen

modified hydroxylated film was formed at 90K, in UHV, by condensing the VTMS

precursor on a titanium substrate, immediately followed by water physisorption,

ammonia physisorption, and electron bombardment, and then annealing to 300K,

resulting in a nitrogen modified hydroxylated amorphous Si:C:H film (a-Si:C:H/OH/N).

XPS measurements indicate Cu deposited by sputter deposition (PVD) at 300 K onto an

a-Si:C:H/N film grows conformally, as indicated by the break of the Cu/Si uptake curve,

similar to, but not as defined as the wetting seen on hydroxyl modified a-Si:C:H films [8].

63

In contrast, conformal growth of Cu deposited on the nitrogen modified hydroxylated film

(a-Si:C:H/OH/N) is indicated by both a break in the Cu/Si uptake curve, and a broad Cu

Auger spectrum from the formation of Cu(I) and Cu(0) species at the interface. These

results are similar to that seen on the hydroxyl modified a-Si:C:H film [8]. As found in

previous studies the unmodified VTMS precursor, physisorbed on titanium and

bombarded by electrons at 90K, and annealed to 300K did not produce charge transfer

or wetting with PVD copper, [4, 8].

3.2 Experiment

These experiments were conducted at temperatures between 90 K and 1200 K,

in an ultra-high vacuum (UHV) system comprised of a main chamber and sputter

deposition chamber, connected by a sample transfer port, enabling transfer without

exposure to the atmosphere, with temperature controlled by liquid nitrogen cooling and

resistive thermocouple heating. The UHV system included a main chamber, with a

base pressure of 6 x 10-10 Torr, and a dual-magnetron sputter deposition chamber, with

a base pressure of 5 x 10-9 Torr [16].

X-ray photoelectron spectroscopy (XPS) spectra were acquired by an non-

monochromated MgKα (1253.6 eV) X-ray source at 15 kV and 300 W with a commercial

hemispherical analyzer operated at 50 eV constant pass energy, in the UHV main

chamber as previously described [16]. XPS spectra were acquired with the analyzer

aligned along the surface normal (normal emission) and at glancing emission from

surface normal. Glancing emission did not produce consistent results, suggesting

64

surface roughness. Therefore the spectra reported here are with the analyzer at normal

emission.

The substrate sample for this study was a 1cm2 area titanium foil (99.97% purity,

0.0076 cm thickness) that was spot welded to titanium leads with a chromel-alumel

thermocouple spot welded to the back for heating and temperature monitoring. Cycles

of Ar+ gas bombardment at a pressure of 1 X 10-5 Torr, alternated with annealing, were

used to clean the sample. Cooling of the sample was done indirectly through the

titanium leads by nitrogen gas that was cooled by liquid nitrogen. VTMS (97% purity),

from a commercial source, was purified by cycles of freeze-thaw pumping, then

delivered ≈ 2 cm from the face of the Ti sample through a stainless steel doser tube

regulated by a manual leak valve. The concentration of VTMS delivered was measured

as gas exposure in Langmuir, L, (1L = 1 x 10-6 Torr s), uncorrected for ion gauge

sensitivity and flux to the surface.

Unmodified VTMS films were produced by condensing 30L of VTMS onto clean

Ti substrates at 90K, immediately followed by electron beam bombardment, to induce

crosslinking [5, 8, 17]. A Kimball Electronics electron gun was used to produce an

electron beam of 500eV energy (3µA emission current) with an average flux at the

sample surface of 1.5 x 10-13 electrons cm-2 s-1 [5, 8]. Nitrogen-modified films were

produced by condensing 30L of VTMS onto clean Ti substrates at 90K, immediately

followed by 30L of NH3, followed by electron beam bombardment. Nitrogen and

hydroxyl-modified films were produced by condensing 30L of VTMS onto clean Ti

substrates at 90K, immediately followed by 30L of H2O, followed by 30L of NH3 and then

followed by electron beam bombardment. All of the films were annealed to 300K in

65

UHV, after electron bombardment, to desorb any remaining condensed species.

Copper was sequentially sputter deposited (PVD) onto each annealed film’s surface, at

300K, using an Advanced Energy MDX 500 magnetron source at 25 watts, while

monitoring the Ar+ pressure with an MKS baratron gauge achieving a rate of ≈ 0.5

Å/minute. After copper deposition the sample was annealed by resistive heating,

through the thermocouple, stepwise from 300K to 1000K. Each step in temperature

was characterized by XPS to detect any resultant interactions between the PVD copper

and the film.

Film composition for both the unmodified and modified films was calculated using

relative atomic concentrations (NA/NB) for carbon, silicon, nitrogen and oxygen, from

XPS core level data with corresponding intensities (I), and appropriate elemental

sensitivity factors (A) for this analyzer [18], using the general equation:

NA/NB = [IAAB]/ [IBAA] (2)

Very thin a-Si:C:H films were produced, which allowed the XPS titanium peak to

be observed after annealing to 300K. The thickness, (d), of the a-Si:C:H films was

calculated as a function of the primary electron flux on the sample, using (ITi), the

intensity of the Ti(2p3/2) core level peak, (λ), and the inelastic mean free path of the

Ti(2p3/2) photoelectron, which is 21.6 Å [16].

ITi = I Ti (0)exp(-d/λ) (3)

The thickness of copper deposited on the modified films was estimated from the

Si(2p) intensity using the intensity ratio of copper to silicon. The intensity ratio of

reference elements, I∞Cu/I∞

Si, the copper to silicon intensity ratio, ICu/ISi, and the inelastic

66

mean free paths for copper (λCu) and silicon (λSi) are known, therefore the thickness

(dCu) can be derived from the equation:

ICu/ISi = (I∞Cu/I∞

Si)[1-exp-dCu/λCu)/exp-dSi/λSi] (4)

Absolute intensities for copper and titanium, calculated using this experimental

system, should have an error in measurement of no more than 10% [5].

3.3 Results

3.3.1 Nitrogen Modified Amorphous-Si:C:H Film Characterization

The average thickness of the unmodified a-Si:C:H cross-linked film produced

from the VTMS precursor, by the process described above, was determined (Eq. 3) to

be 20 Å. As shown in Figure 3.1, the Si(2p) XPS spectrum exhibits a peak at 100.8 eV

attributed to Si-C:H from VTMS [19], and a C(1s) spectra and a peak at 284.5 eV

attributed to C-C and C-H bond types [19], and an insignificant oxygen peak consistent

with previous results [5, 8]. The nitrogen modified silicon:carbon:hydrogen film, (a-

Si:C:H/N), had an average thickness of 27 Å, and the a-Si:C:H/OH/N films average

thickness was 31 Å. Figure 3.2 shows the normal emission spectra for the nitrogen and

nitrogen-hydroxyl modified films. The Si(2p) spectrum for the nitrogen modified film was

fit with a high binding energy peak at 102.2 eV, which was close to the 102.4 eV peak

attributed to Si-N by [20, 21], because of the limited presence of oxygen, and a 100.8

eV peak for Si-C [19]. The C(1s) spectrum included a 284.5 eV peak for C-C and C-Si,

and 285 eV [22] or 285.5 eV for C-N [21] while the nitrogen spectrum consisted of

peaks at 398.1 eV for N-Si and Si-N-O [23], and 399.2 eV for N-C [24, 25]. The

67

nitrogen-hydroxyl modified film was composed of the Si-C 100.8 eV peak and two high

binding energy peaks, 101.9 eV Si-N and 103.8 eV for Si-OH [26] which is in the

Fig. 3.1 Unmodified a-Si:C:H film showing single O(1s), C(1s) and Si(2p) species.

vicinity of 103.7 eV assigned to Si-O [21]. The C(1s) spectrum exhibits a low binding

energy feature at 283 eV, related to a species of carbon possibly bonded to the titanium

substrate, Ti/SiC [19, 27]. In addition, there is a C-N peak at 285.7 [21] and a peak at

284.5 eV for C-C and Si-C, indicating that the oxygen containing species are not

bonded to carbon atoms. The N(1s) spectrum consists of peaks at 398.1 eV for N-Si

and 399.2 eV for N-C [24, 25] and 400.8 eV for N-O [28]. The O(1s) peak (Fig. 3.2) is

fit with two components with FWHM of 2.9 eV and binding energies of 531.3 eV for NOx

and 532.9 eV for Si-OH [8, 29]. Previous work suggested the Si-OH formed in the

-540 -536 -532 -528 -524

-291 -288 -285 -282 -279

-108 -105 -102 -99 -96

C(1s)

Si(2p)

O(1s)

531.3

100.8

284.5

Binding Energy (eV)

XP

S I

nten

sity

(cps

)-540 -536 -532 -528 -524

-291 -288 -285 -282 -279

-108 -105 -102 -99 -96

C(1s)

Si(2p)

O(1s)

531.3

100.8

284.5

Binding Energy (eV)

XP

S I

nten

sity

(cps

)

68

Fig. 3.2 Comparison of XPS spectra (a) nitrogen modified a-SI:C:H film and (b) nitrogen-hydroxyl modified a-Si:C:H film.

hydroxylated a-Si:C:H films as the site for copper wetting [8]. Since there is no hydroxyl

group apparent in the nitrogen modified a-Si-C:H film environment, copper wetting may

be due to electronegative nitrogen groups C-N or Si-N, and the nitrogen-hydroxyl

modified a-Si:C:H films may enable copper wetting at the hydroxyl and nitrogen sites.

C-OH is eliminated as a possibility in the nitrogen-hydroxyl modified film because there

is no characteristic peak near 286 eV in the C(1s) spectra, and in the nitrogen modified

film there is no O (1s) peak near 533 eV, and only small peaks at 532.15 eV for ambient

oxygen contamination [30]. Table 3.1 summarizes the film component binding energies.

-538 -536 -534 -532 -530 -528 0

2000

-405 -402 -399 -396 -3930

1500

Binding Energy (eV)

XPS

Inte

nsity

(b)

-406 - -402 - -398 - -394 0

1500

-106 -104 -102 -100 -98 -96 0

2500

(a)

Si(2p)

C(1s)

O(1s)

N(1s)

-291 -288 -285 -282 -279 0

11000

-290 -288 -286 -284 -282 -2800

13000

-108 -105 -102 -99 -960

3000

-537 -534 -531 -5280

1000

69

Calculation of the C(1s)/Si(2p)total ratio, using appropriate sensitivity factors,

correcting for differences in mean free paths, yielded a C/Si ratio of 3.7:1, for a-

Si:C:H/N, and a C/Si ratio of 3.4:1, for a-Si:C:H/OH/N. These results are similar to the

C/Si ratio of 3.8:1, previously found for unmodified a-Si:C:H [5]. Prior work has found

that chemical and matrix effects, due to differences in kinetic energies, can be corrected

for by dividing by the appropriate mean free path [5]. This treatment was applied to

these results using calculated inelastic mean free path values of for C(1s) = 22.5 Å and

Si(2p) = 25.7 Å and N(1s) = 20.7 Å. The corrected C/Si ratio for a-Si:C:H/N films

increased to ≈ 4.6:1, and the corrected C/Si ratio for a-Si:C:H/OH/N films increased to ≈

4.3:1. Because of the uncertainty of the actual additional chemical and matrix effects

contributed by nitrogen in the a-Si:C:H/N film, and the effects of nitrogen and oxygen in

Si(2p)C(1s)N(1s)O(1s)a-Si:C:H/N

TiO2

Ti/Si-CSi-OHN-OC-NC-CSi-CSi-N

C-NC-CSi-CSi-N

Si(2p)C(1s)N(1s)O(1s)a-Si:C:H/OH/N

102.4398.1100.8284.5

284.5285.5399.2

283530

103.8532.9400.8531.3

288.5399.2284.5

100.8285.7398.1

Table 1 Modified a-Si:C:H Film Binding Energies (eV)

Si(2p)C(1s)N(1s)O(1s)a-Si:C:H/N

TiO2

Ti/Si-CSi-OHN-OC-NC-CSi-CSi-N

C-NC-CSi-CSi-N

Si(2p)C(1s)N(1s)O(1s)a-Si:C:H/OH/N

102.4398.1100.8284.5

284.5285.5399.2

283530

103.8532.9400.8531.3

288.5399.2284.5

100.8285.7398.1

Table 1 Modified a-Si:C:H Film Binding Energies (eV)

70

the a-Si:C:H/OH/N film, a comparison calculation was done using only the a-Si:C:H

Si(2p)(101) peak intensity to calculate the C/Si ratio for the unmodified film that resulted in

a C/Si ratio of 4.7:1. The comparison calculation using only the a-Si:C:H Si(2p)(101)

peak intensity to calculate the C/Si ratio, resulted in a ratio of ≈ 4.8:1 for the a-Si:C:H/N

film, and a C/Si ratio of ≈ 4.6:1 for the a-Si:C:H/OH/N film. Both methods of calculating

the C/Si ratio for a-Si:C:H/N, correlate with the five to one C/Si stoichiometry of VTMS.

The O(1s)total/Si(2p) total atomic ratio for the same a-Si:C:H/N film was calculated to be

0.48:1 and the N(1s)total/Si(2p) total atomic ratio was calculated to be 0.29:1. The

O(1s)total/Si(2p) total atomic ratio for the a-Si:C:H/OH/N film was calculated to be 0.60:1

and the N(1s)total/Si(2p) total atomic ratio for the same a-Si:C:H/OH/N film was calculated

to be 0.31:1. Whereas for unmodified a-Si:C:H the C(1s)/Si(2p) ratio was 4.7:1, and the

O(1s)total/Si(2p) total atomic ratio was 0.11:1. These results are similar to those found for

the hydroxyl modified film in Chapter 2, which found a C/Si ratio of 4.5:1 that correlates

well with the 5:1 VTMS stoichiometry, and a O(1s)total/Si(2p) total atomic ratio calculated

to be 0.44:1 [8].

3.3.2 Copper Deposition by Physical Vapor Deposition (PVD)

XPS was used to characterize the interactions of sputter-deposited Cu with

unmodified and modified film substrates. Cu(2p3/2) core level spectra were used to

determine Cu coverage, and x-ray excited Cu(L3VV) Auger spectra were used to

distinguish Cu(I) vs. Cu(0) oxidation states [7]. The presence of Cu(II) was monitored

by the presence or absence of typical satellite peaks in the Cu(2p3/2) spectra [31, 32],

and was not observed at any point in this experiment (Fig. 3.3). Copper was deposited

71

in the sputter deposition chamber at 300K at sputtering intervals of several minutes and

XPS data were obtained after each deposition.

Fig. 3.3 XPS spectrum showing no Cu(II) satellite peaks upon copper deposition.

No significant changes were observed for C(1s) spectra or Si(2p) spectra

intensities or binding energies after copper deposition on unmodified films (not shown).

In contrast, Cu deposition on the N-modified film results in relative attenuation of the

N(1s) intensity attributed to C-N sites (Fig. 3.4). In the a-Si:C:H/OH/N film the Si-N

102.2 eV peak was attenuated as well as the 103.8 eV Si-OH peak (Fig. 3.5). Copper

growth on the surface of nitrogen modified a-Si:C:H films, depicted in the plot of

Cu(2p3/2)/Si(2p) intensity vs. Cu deposition time (Fig. 3.6), shows a break in the uptake

curve. This break indicates a change in growth mode from conformal to 3-D island

growth similar to Stranski-Krastanov (S-K) growth, whereas unmodified a-Si:C:H films

exhibit consistent linear behavior corresponding to 3-dimensional islanding (Volmer-

-960 -950 -940 -930

0

10000

Binding Energy, eV

XPS

Inte

nsity

72

Weber type growth) [33]. Nitrogen and hydroxyl modified films produce a much more

definite break in the uptake curve for Cu(2p 3/2)/Si(2p) intensity vs. Cu deposition time

Fig. 3.4 XPS spectra for O, C, Si and N species during copper deposition on a-Si:C:H/N film.

shown in Figure 3.7 consistent with conformal growth, and similar to (S-K) growth [33],

albeit the change is much more apparent for the a-Si:C:H/OH/N films.

-405 -402 -399 -396 -3930

1500

-405 -402 -399 -396 -3930

1500

-405 -402 -399 -396 -3930

1500

-405 -402 -399 -396 -3930

1500

-108 -105 -102 -99 -960

3000

-290 -288 -286 -284 -282 -2800

13000

-537 -534 -531 -5280

1000O(1s)

8.5

0.0

6.5

2.5

Binding Energy (eV)

-108 -105 -102 -99 -960

3000

-108 -105 -102 -99 -960

3000

-537 -534 -531 -5280

1000

-537 -534 -531 -5280

1000

-290 -288 -286 -284 -282 -2800

13000-289 -287 -285 -283 -281

0

13000

-537 -534 -531 -5280

1000

XP

S I

nten

sity

(cps

)

Cop

per D

epos

ition

(min

)

N(1s)

C-N

Si-N

-108 -105 -102 -99 -960

3000 Si(2p)

Si-N

Si-C

-290 -288 -286 -284 -282 -2800

13000 C(1s)

C-N

S i-C, C-C

N-OO-O

-405 -402 -399 -396 -3930

1500

-405 -402 -399 -396 -3930

1500

-405 -402 -399 -396 -3930

1500

-405 -402 -399 -396 -3930

1500

-108 -105 -102 -99 -960

3000

-290 -288 -286 -284 -282 -2800

13000

-537 -534 -531 -5280

1000O(1s)

8.5

0.0

6.5

2.5

Binding Energy (eV)

-108 -105 -102 -99 -960

3000

-108 -105 -102 -99 -960

3000

-537 -534 -531 -5280

1000

-537 -534 -531 -5280

1000

-290 -288 -286 -284 -282 -2800

13000-289 -287 -285 -283 -281

0

13000

-537 -534 -531 -5280

1000

XP

S I

nten

sity

(cps

)

Cop

per D

epos

ition

(min

)

N(1s)

C-N

Si-N

-108 -105 -102 -99 -960

3000 Si(2p)

Si-N

Si-C

-290 -288 -286 -284 -282 -2800

13000 C(1s)

C-N

S i-C, C-C

N-OO-O

73

Figure 3.8 chronicles the evolution of Cu(L3VV) spectra during copper deposition

on both a-Si:C:H and a-Si:C:H/N films. For the first minute of copper deposition, ~ 0.5Å,

the Cu(L3VV) Auger spectra showed peaks at 914.9 eV and 918.0 eV, similar to the

formation of both Cu(I) and Cu(0) [34, 35], indicative of conformal growth [8]. With

Fig. 3.5 XPS spectra for O, C, Si and N during copper deposition on a-Si:C:H/OH/N film.

increased copper deposition, Cu films of ~ 1.0Å or more, the Cu(L3VV) peak narrowed

and shifted to 918.8 eV. This was unlike hydroxyl modified a-Si:C:H films, suggesting

weaker charge transfer for a-Si:C:H/N films as compared to Si:C:H/OH films, enabling

-405 -402 -399 -396 -393 -390

0

1400

XPS

Inte

nsity

-536 -533 -530 -5270

8000

-405 -402 -399 -396 -393 -390

0

1400

-405 -402 -399 -396 -393 -390

0

1400

-405 -402 -399 -396 -393 -390

0

1400

-106 -104 -102 -100 -98 -960

2500

-106 -104 -102 -100 -98 -960

2500-106 -104 -102 -100 -98 -96

0

2500-106 -104 -102 -100 -98 -96

0

2500

-536 -533 -530 -5270

8000

-536 -533 -530 -5270

8000

-291 -288 -285 -282 -2790

11000

-291 -288 -285 -282 -2790

11000

-291 -288 -285 -282 -2790

11000

-291 -288 -285 -282 -2790

11000

Binding Energy (eV)

N(1s)Si(2p)O(1s) C(1s)

-536 -533 -530 -527

0

2000

8.5

0.0

6.5

2.5

Cop

per D

epos

ition

(min

)

C-N Si -NSi -N

Si -C

C-N

Si -C, C-C

Ti-CN-O

O-HN-O

Si -OH

Ti-O2

-405 -402 -399 -396 -393 -390

0

1400

XPS

Inte

nsity

-536 -533 -530 -5270

8000

-405 -402 -399 -396 -393 -390

0

1400

-405 -402 -399 -396 -393 -390

0

1400

-405 -402 -399 -396 -393 -390

0

1400

-106 -104 -102 -100 -98 -960

2500

-106 -104 -102 -100 -98 -960

2500-106 -104 -102 -100 -98 -96

0

2500-106 -104 -102 -100 -98 -96

0

2500

-536 -533 -530 -5270

8000

-536 -533 -530 -5270

8000

-291 -288 -285 -282 -2790

11000

-291 -288 -285 -282 -2790

11000

-291 -288 -285 -282 -2790

11000

-291 -288 -285 -282 -2790

11000

Binding Energy (eV)

N(1s)Si(2p)O(1s) C(1s)

-536 -533 -530 -527

0

2000

8.5

0.0

6.5

2.5

Cop

per D

epos

ition

(min

)

C-N Si -NSi -N

Si -C

C-N

Si -C, C-C

Ti-CN-O

O-HN-O

Si -OH

Ti-O2

74

the predominance of copper(0) from the coalescence of Cu adatoms into Cu clusters

[34, 35]. Deposition on the unmodified a-Si:C:H film (Fig. 3.7) resulted in the exclusive

formation of the typical sharp Cu(0) peak initially observed near 916 eV that shifted to

higher binding energies upon additional deposition. Such a shift was seen before for

Fig. 3.6 Cu(2p 3/2)/Si(2p) uptake curve indicating a change in deposition mode as copper is deposited on a-Si:C:H/N films.

Fig. 3.7 Cu(2p 3/2)/Si(2p) uptake curve indicating a change in deposition mode as copper is deposited on a-Si:C:H/OH/N films.

05

10152025303540

0 5 10 15 20 25Deposition Time (min)

Cu/

SiX

PS

Inte

nsity

Rat

io

Cu(I)

Cu(0)

VTMS-N/OH

VTMSa-Si:C:H

a-Si:C:H/OH/N

Break Point

0

2

4

6

8

10

12

14

16

18

20

0.0 2.0 4.0 6.0 8.0 10.0

Deposition Time (min)

Cu/

SiX

PS In

tens

ity R

atio

Cu(I)

Cu(0)

a-Si:C:H/N

a-SI:C:H

Break Point

75

the unmodified a-Si:C:H film [5], and is characteristic of the growth of Cu(0) clusters

[36], due to enhanced screening of Auger final state holes that takes place with

increased copper cluster size [37].

For Cu deposition on the N-modified film, the narrowing of the Auger peak shape

depicted in Figure 3.8, and the increase in kinetic energy that corresponds with the

evolution of the Cu(0) feature, unlike the a-Si:C:H/OH/N film, does not correlate with the

break point in the corresponding Cu(2p3/2)/Si(2p) uptake curve. This is due to weak

charge transfer at the interface of the deposited copper and the modified a-Si:C:H/N

Fig. 3.8 Copper deposited on a-Si:C:H/N films and a-Si:C:H films show only Cu(0) species shifting to higher kinetic energy due to increase in cluster size and final state shielding.

film, which contrasts with the behavior previously observed for unmodified a-Si:C:H

films [8] that produced no change in the uptake curve, as expected for typical Volmer-

Weber growth (3-D islanding) [12]. The increase of kinetic energy of the Auger peak

890 900 910 920 930

20.5

890 900 910 920 930

1.03.56.5

8.5

14.5

a-Si:C:H Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

a-Si:C:H/N Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

0.0

6.5

3.5

2.5

1.0

5.5

8.5

Cu(0)Cu(I) Cu(0)

890 900 910 920 930

20.5

890 900 910 920 930

1.03.56.5

8.5

14.5

a-Si:C:H Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

a-Si:C:H/N Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

0.0

6.5

3.5

2.5

1.0

5.5

8.5

Cu(0)Cu(I) Cu(0)

76

and its narrowing shape is similar to previous findings [5, 8] and are the result of

shielding effects due to Cu(0) cluster formation [34, 35].

In contrast, Cu deposition on the film modified by both nitrogen and hydroxyl

groups (Fig. 3.9) shows a change in Auger peak shape, and change in kinetic energy,

which corresponds with the break point in the Cu(2p3/2)/Si(2p) uptake curve as shown in

Fig. 3.9 a) Copper deposited on a-Si:C:H/OH/N films show Cu(I) and Cu(0) species simultaneously, producing the broad peak shape.

b) Copper deposited on a-Si:C:H films show only Cu(0) species, indicated by narrow peak shape that shifts to higher kinetic energy due to increase in cluster size and final state shielding.

Figure 3.7. For the first four and a half minutes of copper deposition, ~ 2Å, or one

monolayer of Cu on the nitrogen-hydroxyl modified film at 300K (Fig. 3.8), the Cu(L3VV)

Auger spectra showed peaks at 914.9 eV and 918.0 eV, consistent with the formation

of both Cu(I) and Cu(0) [34, 35]. In this case, the behavior is similar to conformal

890 900 910 920 930 890 900 910 920 930

Cu(I) Cu(0)

Kinetic Energy (eV)

Aug

er In

tens

ity

Deposition Time (min)

a-Si:C:H-N/OH Film a-Si:C:H Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

Cu(0)

20.5

14.5

8.5

6.53.51.0

20.5

14.5

8.5

4.52.51.0

6.5

10.512.5

16.518.5

890 900 910 920 930 890 900 910 920 930

Cu(I) Cu(0)

Kinetic Energy (eV)

Aug

er In

tens

ity

Deposition Time (min)

a-Si:C:H-N/OH Film a-Si:C:H Film

Aug

er In

tens

ity

Deposition Time (min)

Kinetic Energy (eV)

Cu(0)

20.5

14.5

8.5

6.53.51.0

20.5

14.5

8.5

4.52.51.0

6.5

10.512.5

16.518.5

77

growth as previously described for hydroxyl-modified a-Si:C:H films [8], with the typical

change in slope of the uptake curve for growth modes similar to S-K growth [33],

followed by the expected increase in kinetic energy and narrowing Auger peak shape

resulting from Cu(0) cluster formation [34, 35].

3.3.3 Annealing Effects

In order to assess the stability of the Cu/substrate interface for the a-

Si:C:H/OH/N film, Cu was deposited on an a-Si:C:H/OH/N film substrate at 300 K,

followed by 15 minute anneals at increasing temperatures in UHV, with corresponding

incremental XPS acquisition at 300K. Results are shown in Figure 3.10 for a copper

film with an initial average thickness of 10 Å prior to annealing. The Cu/Si ratio is

relatively constant until 800K, after which there is a significant decrease in relative Cu

intensity. As previously stated, this decrease in relative copper intensity can been

attributed to either island formation (dewetting) or diffusion of Cu from the surface into

the bulk [4].

A measurement of the relative intensities of the Cu(2p3/2) and Cu(3p) was used to

distinguish between these two possibilities due to the marked differences in the inelastic

electron mean free paths, λ, of the two copper species, 13.9 Å for Cu(2p3/2) and 38.2 Å

for Cu(3p) [12]. Due to its shorter λ, the Cu(2p3/2) peak intensity is preferentially

attenuated relative to the Cu(3p) intensity if copper diffuses into the bulk.

A plot of the Cu(3p)total/Cu(2p3/2)total intensity ratio against temperature (Fig. 3.10)

shows some change below 800 K attributed to dewetting [4], and a significant decrease

78

in Cu(2p3/2) peak intensity and corresponding increase in the Cu(3p)/Cu(2p3/2) ratio [38]

above 800 K, indicating that annealing to temperatures of 800 K or higher induces Cu

diffusion into the nitrogen-hydroxyl modified a-Si-C:H filmas observed for Cu deposited

Fig. 3.10 Copper diffusion is signaled by the rapid increase in the Cu(3p)/Cu(2p) ratio.

on OH-modified a-Si:C:H films [5, 8], and as previously reported for unmodified a-Si:C:H

films deposited on Cu substrates [4].

Annealing above 800 K caused a large decrease in copper auger intensity (Fig.

3.11) and significant changes in the silicon XPS spectra of the a-Si:C:H/OH/N

0

10

20

30

40

0 200 400 600 800 1000 1200

Temperature, K

XP

S In

tens

it y C

u(2 p

)/Si (2

p), c

p s

XP

S In

tens

ity C

u32p

)/Cu(

2 p),

cps

10

20

30

40

0

10

20

30

40

0

Diffusion indicated

79

Fig. 3.11 Copper Auger intensity decreases with anneal from 300 K – 1000 K, and the peak becomes sharper

as it becomes more metallic.

films, suggesting complete breakdown of the films. The dominant Si(2p) 100.8 eV peak

(Si-C:H) at 300K, significantly decreased in intensity after 800K, with the appearance of

two lower energy peaks. The peak at 98.1 eV was attributed to TiSi, with the 99.6 eV

peak attributed to Si-Si bonding. At 1000 K the Si(2p) 98.1 eV (TiSi) peak became the

dominant peak, followed in intensity by the 100.8 eV (Si-C:H) peak with only 10% of its

300K intensity, and a small 99.6 eV (Si-Si) peak.

3.4 Discussion

The evolution of the Cu(0) Auger feature for deposition on the a-Si:C:H/OH/N

film (Fig. 3.9) correlates with the corresponding points in the Cu(2p3/2)/Si(2p) uptake

890 900 910 920 930

0

5000

10000

15000

20000

1000

300

800

900

300 K

800 K

1000 K

Kinetic Energy (eV)

XP

S A

uger

Inte

nsity

(cps

)

Tem

pera

ture

(K)

890 900 910 920 930

0

5000

10000

15000

20000

1000

300

800

900

300 K

800 K

1000 K

Kinetic Energy (eV)

XP

S A

uger

Inte

nsity

(cps

)

Tem

pera

ture

(K)

80

curve, depicted in Figure 3.7. The curve for the a-Si:C:H/OH/N film shows a break that

is characteristic of initial layer by layer (conformal) growth [33] with 3-D islanding, similar

to the growth seen for the OH modified a-Si:C:H film [8]. Conformal growth occurs on

the modified a-Si:C:H/OH/N film, as is indicated by significant attenuation of Si(2p)

Si:C:H intensity from the substrate during copper deposition on the A-Si:C:H/OH/N film.

This correlates with previous research that showed the conformal growth of Cu on OH-

modified a-Si:C:H films [8], whereas only weak initial conformal growth occurs on the

Si:C:H/N film (Figs. 3.6 and 3.8). Evidence for this weak growth is seen in the copper

Auger spectrum that shows both Cu(I) and Cu(0) species for the first minute of

deposition and only Cu(0) after the first minute. In contrast, the unmodified a-Si:C:H

film showed no change in the uptake slope, consistent with solitary Volmer-Weber

growth (3-D islanding) [12, 33], and a corollary sharp Cu Auger peak depicted in Figure

3.8, indicating only Cu(0) coalescence and cluster formation [34, 35].

The thickness of copper deposited on the nitrogen and nitrogen/hydroxyl

modified films was estimated from the Si(2p) intensity using the intensity ratio of copper

to silicon film substrate (Eq. 4). The average thickness of copper at the break point in

the nitrogen modified film’s Cu/Si uptake curve (Fig. 3.8), and for the nitrogen/hydroxyl

modified film’s Cu/Si uptake curve (Fig. 3.9) is ~ 2 Å, comparable to the Cu(I) ion

diameter of 1.92 Å [33] and similar to results for hydroxyl modified films [8]. Since the

initial deposition of Cu is marked by the formation of both Cu(I) and Cu(0), it is difficult to

give a precise determination of surface coverage in the first layer due to the roughness

of the film’s surface. These data estimate that the coverage of the initial Cu adlayer is ~

0.5 slightly less than 1 monolayer. The copper uptake curve data in Figure 3.7 indicate

81

that Cu deposition at 300 K on the a-Si:C:H/OH/N film results in initial layer-by-layer

growth, with the first layer including Cu(I) species along with some Cu(0) species. Upon

completion of the first layer, a second, more metallic Cu layer begins to form similar to

Cu deposition at 300 K on the N-modified a-Si:C:H film that also showed conformal

growth according to the uptake curve. However, charge transfer was not substantial

enough to accomplish copper Auger transitions to form much Cu(I) evidenced by the

domination of Cu(0) in the Auger spectra as shown in Figure 3.8. In contrast, Cu

deposited on the unmodified a-Si:C:H films shows only V-W growth indicating non-

wetting behavior.

Similar transitions from non-wetting to wetting behavior for deposited Cu have

been reported for sapphire(0001) surfaces [7], and for Rh deposited on ordered ultrathin

alumina films [6] and for hydroxyl-modified a-Si:C:H films [8]. In addition, Cu wetting

has also been reported for deposition at 300 K on hydroxylated ultrathin Ta silicate films

on Si(100) and for air-exposed TaSiN substrates [39], and theoretical studies indicate

that Pt nucleation occurs preferentially at surface OH sites on MgO single crystal

surfaces [40]. These results suggest that surface nitridation does not seem to interfere

with surface hydroxylation, enabling hydroxyl-modified films additionally modified by

nitrogen for use as copper adhesion promoters.

Copper diffusion was inhibited by the a-Si:C:H/OH/N film in a manner similar to

unmodified a-Si:C:H film at temperatures below 800K in UHV [4], as shown by a similar

change in the Cu(3p)/Cu(2p3/2) intensity ratio versus temperature curve. Annealing

above 800K caused copper to diffuse into the film and in the additional formation of

Cu(0) at the expense of Cu(I), indicated by the rapid decline in the Cu(3p)/Cu(2p3/2)

82

intensity ratio versus temperature curve. These results are similar to those found for

OH-modified a-Si:C:H films [8]. However, unlike the OH-modifieda-Si:C:H films,

annealing the a-Si:C:H/OH/N film from 500K to 800K resulted in a decrease in Si(100.8)

peak intensity, and increases in Si(102.3 eV) and Si(103.8 eV) peak intensities. This

may be due to uncharacterized reactions within the a-Si:C:H/OH/N film, or between the

film and the small amounts of ambient oxygen. Annealing to 1000K resulted in a large

decrease in intensity of the Si-C 100.8 eV, Si-N 102.3 eV and Si-OH 103.8 eV peaks,

and the production of a large 98.1 eV peak, most likely Ti3SiC2 and TiC 98.4 eV [41],

and a small peak at 99.6 eV, attributed to Si-Si [41]. These peak changes suggest the

total break down of the modified a-Si:C:H/OH/N film.

3.5 Conclusions

This study demonstrates that nitrogen modification of hydroxyl-modified a-Si:C:H

films enables Cu wetting and diffusion inhibition to 800 K in a manner similar to hydroxyl

modified a-Si:C:H films [8], as characterized by x-ray photoelectron spectroscopy

(XPS). In addition, OH/N a-Si:C:H films inhibit copper diffusion to 800 K similar to

hydroxyl-modified a-Si:C:H films, with some film compositional changes at temperatures

above 500K. These findings suggest the use of a-Si:C:H/OH/N films as possible

interlayer dielectrics for the deposition of the initial copper seed layer in interconnect

circuit. If the a-Si:C:H/OH/N film’s dielectric constant is lowered, as in previous

research, to the point that nitrogen modification decreases the dielectric constant and

leakage current [42], these films may be useful as intermetal dielectrics. Additional

83

adaptations of hydroxyl-modification of a-Si:C:H films, substituting other species in the

place of nitrogen, may result in further tailoring the properties of a-Si:C:H/OH films.

This study also demonstrates that nitrogen modification of a-Si:C:H films enables

Cu wetting as compared to unmodified a-Si:C:H films [5, 8], while approximating the

diffusion inhibition properties of unmodified a-Si:C:H films (not shown) as characterized

by x-ray photoelectron spectroscopy (XPS). These findings suggest the use of nitrogen

modified vinyltrimethylsilane derived a-Si:C:H films as a possible interlayer dielectric for

the deposition of the initial copper seed layer in interconnect circuits. If in accordance

with previous research [42], nitrogen incorporation lowers the film’s dielectric constant,

these a-Si:C:H/N films should have a lower dielectric constant than the more polarizable

a-Si:C:H/OH films, suggesting a-Si:C:H/N films as possible intermetal dielectrics.

Changing the nitrogen modification process, by substituting N2 instead of NH3 as the

source of nitrogen, or incorporating other species may result in additional modifications

of a-Si:C:H films. The resulting in a film that can be tailored in various ways to meet

unforeseen requirements associated with the continued decrease in feature sizes

predicted by Moore’s Law [43].

84

3.6 Chapter References

[1] Liu, P.-T., Chan, T.-C., Yang, Y.-L. and Cheng, Y.-F. S., S.M., IEEE Transactions

on Electron Devices 47 (2000) 1733 -1739.

[2] Vogt, M., Kachel, M., Plötner, M. and Drescher, K., Microelectronic Engineering

37/38 (1997) 181-187.

[3] McBrayer, J. D., Swanson, R. M. and Sigmon, Y. W., J. Electrochem. Soc. 133

(1986) 1242-6.

[4] Chen, L. and Kelber, J. A., J. Vac. Sci. Technol. A 17 (1999) 1968.

[5] Tong, J., Martini, D., Magtoto, N., Pritchett, M. and Kelber, J., Appl. Surf. Sci. 187

(2002) 256-264.

[6] Libuda, J., Frank, M., Sandell, A., Andersson, S., Bruhwiler, P. A., Baumer, M.,

Martenson, N. and Freund, H.-J., Surf. Sci. 384 (1997) 106.

[7] Niu, C., Shepherd, K., Martini, D., Tong, J., Kelber, J. A., Jennison, D. R. and

Bogicevic, A., Surf. Sci. 465 (2000) 163-176.

[8] Pritchett, M., Magtoto, N., Tong, J., Kelber, J. A. and Zhao, X., Thin Solid Films

440 (2003) 100-108.

[9] Lee, S. G., Kin, Y. J., Lee, S. P., Oh, H.-S., Lee, S. J., Kim, M., Kim, I.-G., Kim,

J.-H., Shin, H.-J., Hong, J.-G., Lee, H.-D. and Kang, H.-K., Jpn. J. Appl. Phys. 40

(2001) 2663-2668.

[10] Lanckmans, F., Gray, W. D., Brijs, B. and Maex, K., Microelectron. Eng. 55

(2001) 329-335.

85

[11] Deplancke, M. P., Powers, J. M., Vandentop, G. J., Salmeron, M. and Somorjai,

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88

CHAPTER 4

RESEARCH SUMMARY

Semiconductor device features continue to decrease in size while increasing in

density in response to Moore’s Law resulting exposing the physical, thermal and

electrical limitations of aluminum and silicon dioxide requiring the development of

replacement metal and dielectric systems.

Copper is the obvious choice to replace aluminum due to its superior

electromigration, conductivity, resistivity and thermal properties. However Cu diffuses

into silicon dioxide under normal operating conditions, requiring the use of barrier

materials that inhibit Cu diffusion into SiO2. Barrier layers add to insulator thickness and

delay time, therefore instead of a separate barrier layer. Therefore, another option is to

replace SiO2 with dielectrics that copper can grow on conformally that also inhibit

copper diffusion.

Various amorphous silicon carbon hydrogen (a-Si:C:H) films have been

suggested as possible diffusion barriers and replacements for SiO2 as interlevel

dielectrics (ILD), or intermetal dielectrics (IMD), due to their lower dielectric constants

[1], acceptable thermal conductivity [2] and the ability of copper to grow conformally on

them [3].

This study condensed Vinyltrimethylsilane (VTMS) precursor on titanium

substrates at 90 K and bombarded the physisorbed VTMS with electron beams to

induce crosslinking and formpolymerized a-Si:C:H films. The film was modified by

89

condensing water, ammonia, or water and ammonia, on the condensed VTMS before

electron bombardment, resulting in hydroxyl modified (a-Si:C:H/OH), nitrogen modified

a-Si:C:H/N, and nitrogen and hydroxyl modified (a-Si:C:H/OH/N) polymerized films. X-

ray Photoelectron Spectroscopy (XPS) was used to analyze the films before and after

copper metallization by physical vapor deposition (PVD) to characterize any interactions

between the film and copper species that indicate the wetting behavior of copper, and

XPS was used after annealing to determine the onset of copper diffusion.

XPS results detailed in Chapter 2 show that a-Si:C:H/OH films are composed

primarily of C-C, Si-C, and Si-OH, bonded species that develop strong interactions

enabling the conformal growth of Cu (I) and Cu (0) species, as indicated by Auger

spectra, for a monolayer of copper deposition, that resists thermally-stimulated copper

diffusion at temperatures up to 800K, and complete film dissociation at ~ 1000K.

Findings for nitrogen modified a-Si:C:H films, outlined in Chapter 3, show XPS

spectra dominated by C-C, Si-C and Si-N, bond types for both a-Si:C:H/OH/N and a-

Si:C:H/N films, and Si-OH only present in a-Si:C:H/OH/N films. Copper grows

conformally on a-Si:C:H/OH/N films for the first monolayer of copper deposition when,

according to Auger spectra, both Cu (I) and Cu (0) species are present, and these films

resisted thermally-stimulated copper diffusion at temperatures up to 800K, similar to a-

Si:C:H/OH films. Before these films dissociate at 1000 K, they interact with titanium

forming Ti3SiC2 and TiC species at temperatures above 800K. a-Si:C:H//N films are

composed of C-C, Si-C and Si-N bonds with little oxygen present. Copper grows

conformally on these films according to Cu/Si intensity ratio data. However, the Cu

Auger spectra does not show the presence of Cu (I) and Cu (0) species, suggesting the

90

interface interactions produce only weak charge transfer that lasts for less than 0.5

monolayers of copper deposition before 3-D islanding dominates.

This study demonstrates thatit is possible to modify a-Si:C:H films so that copper

will grow conformally on them, while preserving their ability to wet titanium and inhibit

copperdiffusion, and indicates that the main component responsible for the conformal

growth of copper is the hydroxyl species, not nitrogen.

91

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