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THESIS

Zoc‘

This is to certify that the

dissertation entitled
Diffusion Barriers for Silicon Carbide Particle

Reinforcements by Ion-Beam Assisted Deposition:
Effects On Interphase Stability in SiCp/beta-NiAl
and Sicp/gamma-N13A1 Composites

presented by

Zhiwéi Cai

has been accepted towards fulﬁllment
of the requirements for

Materials Science and
Engineering

 

 

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11m moan-peso.“

 

DIFFUSION BARRIERS FOR SILICON CARBIDE PARTICLE REINFORCEMENT S
BY ION-BEAM ASSISTED DEPOSITION: EFFECTS ON INTERPHASE STABILITY
IN SiCp/B-NiAl AND SiCpl'y-Ni3Al COMPOSITES

By

Zhiwei Cai

A DISSERTATION

Submitted to
Michigan State University
in partial fulﬁllment of the requirements
for the degree of

DOCTOR OF PHILOSOPHY

Department of Materials Science and Mechanics

2000

ABSTRACT

DIFFUSION BARRIERS FOR SILICON CARBIDE PARTICLE REINFORCEMENTS
BY ION BEAM ASSISTED DEPOSITION: EFFECTS ON INTERPHASE STABILITY
IN SiCp/B-NiAl AND SiCp/y-Ni3A1 COMPOSITES

By

Zhiwei Cai

In this study, aluminum nitride and aluminum oxide ﬁlms were used as diffusion
barriers for SiC particles that were consolidated with B-NiAl and y- Ni3Al matrices at
temperatures of 1673 K and 1373 K, respectively. The focus of this study was to
understand factors inﬂuencing the effectiveness of the diffusion barriers during the

consolidation processes of the two composite systems.

The barrier ﬁlms were deposited on SiC particles by ion-beam assisted vacuum
evaporation during which the SiC particles were radiantly heated and acoustically
levitated. The nitride ﬁlm formed reactively on SiC particles, and consisted of 95%
aluminum nitride (balanced with aluminum nitrate and oxide). The oxygen content in the
nitride ﬁlm was a result of the impingement of residual oxygen and water molecules in

the deposition environment.

A voided globular structure of ﬁne-grained clusters was found in a nitride ﬁlm
deposited on SiC particles at 593 K, which was attributed to the levitation of the particles
and the deposition temperature. Nitride ﬁlms deposited at a higher temperature of 793 K
consisted of a ﬁne-grained dense structure with few voids. The oxide ﬁlm deposited at

room temperature had a ﬁne-grained dense structure with some globular features.

This study found that ﬁlm material affected ﬁlm’s ability of retaining integrity
during compositing process, which was important for the success of the barrier ﬁlms.
Annealed at 1673 K, grains in a nitride ﬁlm (deposited at 793 K) coalesced to an average
size of 0.5 pm that was comparable to the ﬁlm thickness. Grain boundaries in the ﬁlm
were widened by the pore agglomeration, resulting in micro-cracks. The oxide ﬁlm
exhibited a similar phenomenon of uninhibited grain growth and micro-crack formation
at a lower temperature of 1273 K. Both ﬁlms failed to be an effective barrier in SiC/[3-

NiAl composite during the compositing process at 1673 K.

This study showed the inﬂuence of ﬁlm structure on grain growth in ﬁlms at high
temperatures. With a voided globular structure, the nitride ﬁlm deposited at 593 K
experienced a grain growth that was primarily restricted within each individual grain
clusters at elevated temperatures. Solid-state sintering at 1673 K densiﬁed the nitride
ﬁlm because of the ﬁlm’s originally voided structure. After annealing at 1673 K for 4
hours, an average 0.5 pm thick nitride ﬁlm deposited at 593 K retained ﬁlm integrity on

SiC particles, and succeeded as a diffusion barrier in the SiC/B-NiAl composite.

This study indicated that the chemical reaction potential in a composite system
was a critical factor in determining the effectiveness of diffusion barriers. SiC/y-Ni3Al
composite had a higher chemical reaction potential at 1373 K, comparing to SiC/B-NiAl
system at 1673 K. Although being effective as a diffusion barrier in the SiC/B-NiAl
composite at 1673 K, the same 0.5 pm thick nitride ﬁlm (deposited at 593 K) did not
protect the SiC particles in the SiC/y-Ni3Al composite at 1373 K.

A cursory study performed in this research indicated that Ni has low diffusivities
in aluminum nitride lattice. From a y—Ni3Al source, the upper—bound lattice diffusivity
was estimated to be 2.2x 10‘l4 cmZ/sec at 1473 K. From a B-NiAl source, the upper-

bound lattice diffusivity at 1673 K was estimated to be 5.8x 10'M cm2/sec.

I dedicate this dissertation to my wife, Fang Li.

And to the memories of my late grandmother.

iv

ACKNOWLEDGEMENTS

I wish to express my sincere appreciation to Dr. D. S. Grummon, my advisor, for
his generous support, patient guidance, creative ideas and scientiﬁc expertise. I would
also like to thank my committee members, Dr. M. A. Crimp, Dr. T. R. Bieler, and Dr. K.
L. Klomparens for their helpful discussions and valuable criticisms of the dissertation.
Also, the generosity of the Composite Materials and Structures Center for letting me use

their equipment is appreciated.

I would like to give a heartfelt thanks to my dear wife, Fang Li, for her love. It
would be impossible for me to ﬁnish this dissertation without her enormous patience and
encouragement throughout the entire unbelievable length of time to complete this

dissertation.

I am deeply indebted to our parents for their caring, support, and wisdom. I am
particularly grateful for our parents’ sacriﬁces in their senior years helping my wife and I

to raise our young family in an environment so alien to them.

I would also like to thank my two little angels, Phillip and Louis, for all the joy

they bring into my life.

CHAPTER 1.

CHART ER 2.

TABLE OF CONTENTS

LIST OF TABLES .................................................................................................

LIST OF FIGURES ................................................................................................

l. Diffusion in Thin Films .........................................................................

General Aspects
Grain Boundary Diffusion

2. Interaction between SiC and Nickel Aluminides ..................................

3. Diffusion Barriers ..................................................................................

Principles of Diffusion Barrier
Candidates for Diffusion Barriers
A1203 as a Diffusion Barrier in SiC/y -Ni3Al System

4. Comparison of Aluminum Oxide and Aluminum Nitride ....................

5. Vacuum Evaporation and Ion-Beam Assisted Deposition ....................

PVD Techniques for Thin Films
Vacuum Evaporation

Reactive IBAD for AlN Films
Introduction
The Ion Source
Reactive [BAD
Aluminum Nitride Film Deposition with Reactive [BAD

Thin Film Deposition on Particulate Substrates

6. Zone Structures of Thin Films from PVD .............................................

vi

INTRODUCTION .......................................................................

BACKGROUND REVIEW ........................................................

...X

...xi

1

...4

...4

4
5

11

13
13
14
16

20

...20

7. Grain Growth and Solid—state Sintering in Ceramics ............................... 22

CHAPTER 3. EXPERIMENTAL PROCEDURES .............................................. 24
1. Materials ................................................................................................... 24

2. Particle Levitation and Heating System ................................................... 25
System Overview 25

Levitation System 26

Particle Holder Assembly
Vibration Generator System

Connector Assembly
Heating Zone 28
Performance of the System 29
3. Deposition System and Procedures .......................................................... 30
IBAD System 30
Aluminum Nitride Film Depositions 31
Aluminum Oxide Film Depositions 33
Deposition Conditions on Particulate Substrates 33
4. Film Characterization ............................................................................... 35
Film Structure and Thickness 35
Composition Study 37
5. Thin Film Therrnomechanical Behavior: Annealing Tests ...................... 38
Annealing at 1273 K in Air 38
Annealing of Aluminum Nitride Films in Vacuum 39
6. Ni Diffusion in Aluminum Nitride, a Cursory Study ............................... 40
Diffusion in Bulk Aluminum Nitride 40
Ni Penetrating Aluminum Nitride Films 41
7. Effectiveness of Diffusion Barriers .......................................................... 41

vii

CHAPTER 4. RESULTS AND DISCUSSION .................................................... 43

1. Film Characterization ............................................................................... 43
Film Thickness and Deposition Rate on SiC Particles 43
Film Structure and Morphology 46

Aluminum Nitride Film Deposited at 593 K

Aluminum Nitride Film Deposited at 793 K

Aluminum Oxide Film Deposited at Room Temperature

Discussion on the Globular Structure in Nitride Films
Composition of Aluminum Oxide Films 49
Composition and Constituents in Aluminum Nitride Films 50

Nitride Film Deposited with an i/a Ratio of 0.42

Nitride Film Deposited with an i/a Ratio of 0.60

Origin of Oxygen inside Nitride Films

Discussion of Nitride Film Deposition Process on Planar

Substrates
Discussion of Composition of Nitride Films on SiC Particles
Summary 60
2. Thin Film Thermomechanical Behavior: Annealing Tests ...................... 61
Annealing at 1273 K in Air: Nitride Films on SiC Particles 61
Annealing at 1273 K in Air: Oxide Films on SiC Particles 62
Summary of Annealing Tests in Air 63
Annealing in Vacuum: Nitride Films Deposited at 793 K 63

Films Annealed at 13 73 K for One Hour
Films Annealed at 1573 K for Four Hours
Films Annealed at 1673 K for Four Hours
Annealing in Vacuum: Nitride Films Deposited at 593 K 65
Films Annealed at 1 5 73 K for Four Hours
Films Annealed at 1673 K for Four Hours

Discussion of Nitride Film Annealing in Vacuum 67
3. Ni Diffusion in Aluminum Nitride, a Cursory Study ............................... 69
Ni Diffusion at the Interface of Aluminum Nitride and
Nickel Aluminide Plates 69
Ni Diffusional Penetration in a Deposited Aluminum
Nitride Films 70
Summary and Discussion 71

4. Effectiveness of Diffusion Barriers in Two Composites: SiC
Particle Reinforced B-NiAl and y-Ni3Al Composites ............................... 72

The Control Composite Samples 73

viii

Aluminum Nitride Films Deposited at 593 K as Diffusion Barrier 74
In ﬂ-NiAl Matrix Composite
In y—Ni 3A1 Matrix Composite

Discussions
Aluminum Nitride Films Deposited at 793 K as Diffusion Barrier 76
In ,B-NiAl Matrix Composite
In y-Ni 3A1 Matrix Composite
Discussions
Aluminum Oxide Films as Diffusion Barrier 77
Summary 78
CHAPTER 5. CONCLUSIONS ............................................................................ 79
APPENDIX 1. PRINCIPLES OF THE XPS TECHNIQUE .................................... 84

APPENDIX 11. CALCULATION OF N ITRIDE FILM MOLECULAR

CONCENTRATION FROM XPS RESULTS ............................... 86

Case 1. Nitride Film with an i/a Ratio of 0.42 on Planar Substrate ............. 86
Interior Region 86

Surface Region 87

Case 2. Nitride Film with an i/a Ratio of 0.60 on Planar Substrate ............. 89
Interior Region 89

Surface Region 90

Case 3. Nitride Film Deposited on SiC Particles ......................................... 91
General Assumptions 91

Film Deposited at 593 K with an i/a Ratio of 0.56 92

Film Deposited at 793 K with an i/a Ratio of 0.53 94
BIBLIOGRAPHY ..................................................................................................... 98

ix

Table 3.1.

Table 4.1.
Table 4.2.

Table 4.3.

Table 4.4.

Table 4.5.

Table A.

LIST OF TABLES

Deposition Rates and i/a Ratios of AlN Films Deposited on Planar
Substrates. ............................................................................................... 32
Measured Thickness of As-Deposited Films on SiC Particles ................ 44
Composition of the Nitride Film Deposited with an i/a = 0.42, at a
Region 40 nm below the Film Surface. ................................................... 51
Composition of the Nitride Film Deposited with and i/a = 0.42, at the
Surface Region. ....................................................................................... 52
Composition of the Nitride Film Deposited with an i/a = 0.60, at the
Surface Region. ....................................................................................... 53
Composition of the Nitride Film Deposited with an i/a = 0.60, at a
Region 20 mm below the Film Surface .................................................... 54

Standard XPS Binding Energy for Al, N, and O [Moulder, 1992;
Taylor, 1981]. .......................................................................................... 85

Figure 2.1.

Figure 2.2.

Figure 2.3.

Figure 2.4.

Figure 2.5.

Figure 2.6.

Figure 2.7.

Figure 2.8.

Figure 2.9.

Figure 3.1.
Figure 3.2.

Figure 3.3.

Figure 3.4.
Figure 4.1.

Figure 4.2.

Figure 4.3.

LIST OF FIGURES

Illustration of Temperature Regimes of Dominant Diffusion

Mechanisms ....................................................................................... 103
Regimes of Dominant Diffusion Mechanism in FCC Metal

Films [after Ballufﬁ, 1975]. .............................................................. 103
A Comparison of Concentrational Proﬁles between the Film

and Bulk Models [after Gilmer, 1976a]. ........................................... 104
Schematic Presentation of Reaction Zone between SiC and

y-Ni3Al at 1273 K for 23 Hours [Chou, 1991b]. ............................... 105
Schematic Presentation of Reaction Zone between SiC and

B-NiAl at 1573 K for 20 Hours [Chou, 1991a]. ................................ 105
A Comparison of Linear Thermal Expansion of AlN, SiC,

and A1203 [after Carborumdum Co., 1992]. ..................................... 106
A Comparison of Thermal Conductivity of AlN, SiC, and

A1203 [after Carborumdum Co., 1992]. ............................................ 106
Schematic Conﬁguration for Ion-Beam-Assisted Electron-

Beam Evaporation. ............................................................................ 107
Schema of Zone Structures in Evaporated Thin Films [after

Movchan, 1969]. ............................................................................... 107
A SEM Photo of an As-Received SiC Particle. ................................ 108
Conﬁguration of an IBAD System for Deposition on

Particulate Substrates. ....................................................................... 109
Illustration of the Arrangement of the Heater Compartment

for the Particle Levitation System. .................................................... 110
Conﬁguration for Hot-Pressing. ........................................................ l l 1

Cross-section of Aluminum Nitride Film on a SiC Particle,
Deposited at 793 K. ........................................................................... 112

Aluminum Nitride Film on a Planar Substrate, Deposited at
593 K. ................................................................................................ 113

Aluminum Nitride Film on a SiC Particle, after a 9.75-hour
Deposition at 593 K ........................................................................... 114

xi

Figure 4.4.

Figure 4.5.

Figure 4.6.

Figure 4.7.

Figure 4.8.

Figure 4.9.

Figure 4.10.

Figure 4.11.

Figure 4.12.

Figure 4.13.

Figure 4.14.

Figure 4.15.

Figure 4.16.

Figure 4.17.

Figure 4.18.

Figure 4.19.

Figure 4.20.

Aluminum Nitride Film on a SiC Particle, after a 25-hour
Deposition at 793 K ........................................................................... 115

Aluminum Oxide Film on SiC Particles, after a 6-hour
Deposition at Room Temperature. .................................................... 116

Aluminum Nitride Film on a SiC Particle, after a 2-hour
Deposition at 593 K ........................................................................... 117

XPS Peaks from an Aluminum Oxide Film Deposited at
Room Temperature, after Curve-ﬁt ................................................... 1 18

XPS Concentrational Proﬁle of an Aluminum Nitride Film
Deposited at 793 K with an i/a of 0.42. ............................................ 119

XPS Peaks from an Aluminum Nitride Film Deposited at
793 K with an i/a of 0.42, 40 mm below the Original Film
Surface. .............................................................................................. 120

XPS Peaks from an Aluminum Nitride Film Deposited at
793 K with an i/a of 0.42, at Surface Region. ................................... 121

The XPS Oxygen Peak from an Aluminum Nitride Film
Deposited at 593 K with an i/a of 0.60, Surface Region ................... 122

The XPS Oxygen Peak from an Aluminum Nitride Film
Deposited at 593 K with an i/a of 0.60, 20 nm below the
Original Film Surface ........................................................................ 122

Aluminum Nitride Film Deposited on SiC Particles at 793 K,
after Air-cooling from 1273 K to Room Temperature. ..................... 123

Aluminum Nitride Film Deposited on SiC Particle at 593 K,
after Quenching in Liquid Nitrogen (77K) from 1273 K. ................. 124

Aluminum Oxide Film on a SiC Particle, after Air-cooling
from 1273 K to Room Temperature. ................................................. 125

Aluminum Nitride Film Deposited on SiC Particles at 793 K,
after Annealing at 1373 K in Vacuum for 1 Hour ............................. 126

Aluminum Nitride Film Deposited on SiC Particles at 793 K,
after Annealing at 1573 K in Vacuum for 4 Hours. .......................... 127

Aluminum Nitride Film Deposited on SiC Particles at 793 K,
after Annealing at 1673 K in Vacuum for 4 Hour ............................. 128

Aluminum Nitride Film Deposited on SiC Particles at 593 K,
after Annealing at 1573 K in Vaccum for 4 Hours. .......................... 130

Aluminum Nitride Film Deposited on SiC Particles at 593 K,
after Annealing at 1673 K in Vaccum for 4 Hours. .......................... 131

xii

Figure 4.21.

Figure 4.22.

Figure 4.23.

Figure 4.24.

Figure 4.25.

Figure 4.26.

Figure 4.27.

Figure 4.28.

Figure 4.29.

Figure 4.30.

Ni Concentrational Proﬁle at Interfaces in Diffusion
Couples, by EDX ............................................................................... 132

XPS Proﬁle of Ni Diffusing in an Aluminum Nitride Film
from a y-Ni3Al Source after a 1373 K/30 Minutes Annealing. ......... 133

Reaction Zone in a SiC Particle Reinforced B-NiAl Matrix
Composite with Uncoated SiC Particles. .......................................... 134

Reaction Zone in a SiC Particle Reinforced Y-NI3AI Matrix
Composite with Uncoated SiC Particles. .......................................... 135

Aluminum Nitride Film Deposited at 593 K as a Diffusion
Barrier in a SiC Particle Reinforced B-NiAl Matrix Composite ....... 136

Aluminum Nitride Film Deposited at 593 K as a Diffusion
Barrier in a SiC Particle Reinforced B-NiAl Matrix
Composite, Film Damaged during Hot-Press Process. ..................... 138

Reaction Zone in a SiC Particle Reinforced y-Ni3Al Matrix
Composite with Aluminum Nitride Film Deposited at 593 K
as a Diffusion Barrier. ....................................................................... 140

Reaction Zone in a SiC Particle Reinforced B-NiAl Matrix
Composite with Aluminum Nitride Film Deposited at 793 K
as a Diffusion Barrier. ....................................................................... 141

Reaction Zone in SiC Particle Reinforced y-Ni3Al Matrix
Composites with Aluminum Nitride Film Deposited at 793 K
as a Diffusion Barrier. ....................................................................... 142

Reaction Zone in SiC Particle Reinforced B-NiAl Matrix
and y-Ni3Al Matrix Composites with Aluminum Oxide
Film as a Diffusion Barrier ................................................................ 143

xiii

CHAPTER 1. INTRODUCTION

Ceramic-reinforced interrnetallic matrix composites are of interest for elevated-
temperature applications because of their better theoretical high-temperature properties as
compared with the materials currently in service [Brindley, 1987; Metcalfe, 1974;
Schoutens, 1982]. However, many of these composite systems suffer from inter-
diffusion and/or interfacial reactions during composite consolidation and application at

high temperatures, which in turn degrades the performance of these composites.

Silicon carbide (SiC) reinforced nickel aluminide-based matrix composites
(SiC/B-NiAl and SiC/y-NigAl systems) are examples of high-temperature composites that
have both potentials and problems. Although attractive from the standpoint of oxidation
resistance and good theoretical mechanical properties at temperatures above 1273 K,
these composites suffer from strong and debilitating interfacial chemical reactions of the
carbide reinforcement with the nickel-bearing matrices. Because of the potential beneﬁts
of these composites as the next generation of high-temperature structural materials,
studies have been done to understand the problems involved and their possible Solutions

[Chou, 1991a; Nieh, 1989; Pope, 1987; Yang, 1989 a&b].

SiC reacts strongly with nickel-bearing intermetallics at temperatures below the
expected composite processing and service temperatures, producing graphite and ternary
compounds. The reaction of SiC with B-NiAl has a modest rate and occurs at
temperatures above 1573 K [Chou, 1991a], while y—Ni3Al reacts with SiC more strongly
even at a lower temperature of 1073 K [Nieh, 1989; Yang, 1989b]. The interface zones
formed have inherently poor mechanical properties. A continuous supply of free Ni

atoms into the SiC region is the controlling factor in these reactions [Chou, 1991 a&b].

For this reason, concept of diffusion barrier against Ni diffusing into SiC was suggested

for improving the viability of these composites.

Aluminum oxide (A1203) thin ﬁlm has been studied as a diffusion barrier
candidate in SiC/y-Ni3Al composite system. The results were not consistent, though.
N ieh [1989] reported that an A1203 ﬁlm formed on the surface of a N i3Al-based alloy
plate demonstrated the ability to function as a diffusion barrier between a planar SiC and
the Ni3Al plate]. However, A1203 ﬁlms deposited on SiC particles by reactive sputtering
consistently failed to protect the SiC particles from the reactions at the consolidation

temperature of 1373 K for y-Ni3Al matrix composites [Cai, 1992].

Because of the inconsistent results of A1203 as a diffusion barrier in SiC/y-Ni3Al
composite system, this study evaluated several factors that inﬂuenced the effectiveness of
diffusion barriers. These factors included barrier ﬁlm materials, barrier ﬁlm deposition
conditions and resultant ﬁlm structures, barrier ﬁlm thermomechanical behavior at

elevated temperatures, and chemical potentials in different composite systems.

In this study, SiC particle reinforced B-NiAl and y-Ni3Al matrix composites were
selected for different levels of chemical potential in the composite systems. Aluminum
nitride (AlN) was selected as a diffusion barrier material for SiC at elevated temperatures
because AlN has more attractive thermomechanical properties (chieﬂy a more favorable
thermal expansion coefﬁcient) than A1203 [Poluboyarinov, 1980; Shaffer, 1964;

Sheppard, 1990]. Aluminum oxide barriers were studied as a basis for comparison.

Both aluminum nitride and oxide ﬁlms were deposited on SiC particles using an

ion-beam assisted deposition (IBAD) technique. A novel particle levitation and heating

 

1 In Nieh's experiment, a Ni3Al plate was pre-oxidized to form a 0.13 pm thick A1203 layer on its surface.
This oxide layer successfully prevented the inter-diffusion and reactions between the aluminide plate and
a SiC plate at 1373 K.

system was designed to deposit ﬁlms uniformly on the SiC particulate substrates. The
nitride ﬁlms were deposited reactively by using a nitrogen ion beam to react with
evaporated aluminum atoms at the substrate surfaces. The composition of the nitride ﬁlm

was studied for evaluating the ability of such a reactive deposition process.

Two factors, the as-deposited ﬁlm structure and ﬁlms’ thermomechanical
behavior during compositing processes, were particularly interested in this study. The
two factors were not only interrelated, but also underlay the viability of the ﬁlms as an
effective diffusion barrier. The effectiveness of the diffusion barriers was evaluated in
model composites under high-temperature consolidation conditions after combining all

the inﬂuencing factors analyzed in this study.

Because of the limited information on Ni diffusing in aluminum nitride, a cursory
investigation was conducted to ensure that Ni has low lattice diffusivity in aluminum
nitride at the temperatures applied in this research. Only the extent of Ni diffusing in

aluminum nitride was of interest in this investigation.

Chapter 2 presents a review of the related ﬁelds of diffusion, the interaction
between SiC and nickel aluminides, the concept of diffusion barriers, a property
comparison between A1203 and AlN, vacuum evaporation and the IBAD processes, the
resultant ﬁlm structures, and grain growth in ceramics. The remaining chapters discuss

the experimental procedures, results, and conclusions drawn from the study.

CHAPTER 2. BACKGROUND REVIEW

This chapter reviews ﬁelds related to this research, including diffusion, diffusion
barriers and barrier materials, methods of thin ﬁlm deposition, and thermomechanical
behaviors in thin ﬁlms during high-temperature treatment. Much of the discussion is

focused on the aspects that directly apply to this study.

1. Diffusion in Thin Films
General Aspects

Diffusion is a time-dependent, thermally activated process that is driven by
gradients in chemical potential, and rate-limited by atomic mobility in a material system.
The rate factor for a diffusant in a material system, the diffusivity (D), can be expressed
in terms of a frequency factor (0,), temperature (T) and an activation energy (Q):

D = Do exp(—-RQ7)
Here, R is the molar gas constant. Both Q and Do depend on the matrix as well as the
diffusant. The distance of diffusion is proportional to 25 , where t is the appropriate

diffusion duration [Gupta, 1988; Shewmon, 1963].

Diffusion in solids involves atomic exchange with lattice imperfections. The
concentration of defects has a profound impact on diffusivity. In ionic compounds with
negligible deviation from their stoichiometries, such as in A1203, the diffusivities are
often small because of low point defect concentrations [Gupta, 1988; Kingery, 1976].

Compared with lattice-bound atoms, less tightly bound atoms at non-lattice sites, such as

grain boundaries, dislocation cores, surfaces, and interfaces, are more mobile. Diffusion
through these sites requires less activation energy than through lattice sites, resulting in
higher diffusivities at low temperatures (so-callfast-diﬁusion) [0hring, 1991]. At high

temperatures, lattice diffusion generally dominates mass transport in materials.

Figure 2.1 illustrates the corresponding dominant temperature regimes for lattice
diffusion and grain-boundary diffusion. In most materials, grain-boundary diffusion
dominates at temperatures below 0.3 Tm, where Tm is the melting point of the material
[0hring, 1991]. The available number of fast-diffusion site inﬂuences the temperature
division between the dominant diffusion mechanisms. Higher defect concentration

pushes the division to a higher temperature.

Compared with bulk materials, thin ﬁlms from physical vapor deposition have
higher density of defects such as vacancies, voids, grain boundaries and dislocations,
which make the fast-diffusion important even at relatively high temperatures. Over a
wide range of temperatures, diffusion in a thin ﬁlm is a combined result of the lattice
diffusion and fast-diffusion mechanisms. Among the latter, grain-boundary diffusion is
the most important because of the small grain sizes of these materials. Figure 2.2
illustrates the regimes of dominant diffusion mechanisms in FCC metal ﬁlms at various
temperatures [after Ballufﬁ, 1975]. For typical metal ﬁlms with a grain size of 1 um or

less, grain-boundary diffusion dominates at all practical temperatures.

Grain Boundary Diffusion

Grain-boundary diffusion can seldom be fully de-coupled from lattice diffusion
since the diffusing species may leak from grain boundaries into the adjoining lattice.

Three types of kinetic behavior, A-, B—, and C-kinetics, are used to describe this lateral

leakage [Gupta, 1978]. Mathematical models have been established to solve the diffusant

concentration for each kinetic regime.

A-kinetics describes the situation where the diffusant has high diffusivity in the
matrix lattice. Extensive lattice diffusion causes the ﬁelds of diffusing species from
adjoining grains to overlap. In C-kinetics, on the other extreme, lattice diffusion is
considered negligible and signiﬁcant atomic transport occurs only within the grain
boundaries. C-kinetics requires very low lattice diffusivity in the diffusion system, and

relatively short diffusion duration [Gupta, 1978]1.

B-kinetics deals with an intermediate situation, where one boundary is assumed to
be isolated from the other and where the lateral ﬂux of the diffusant approaches zero at
large distances from the grain boundary. The diffusant concentration in the lattice is the
sum of the contribution from lattice diffusion and material leakage from the grain
boundaries. Whipple [1954] has given a solution for B-kinetic diffusion in a bulk

material under a constant diffusant source.

The concentration proﬁle for B-kinetics diffusion has two distinguishable regions.
A high concentration region is dominated by the lattice diffusion, with a large gradient
and a shallow penetration depth near the source interface. The followed second region is
of a low level and less gradient of diffusant concentration, which arises from grain-
boundary diffusion [Gilmer, 1976a]. The diffusant concentration from a grain boundary
diffusion is proportional to ya,5 , where y is the diffusion distance [LeClaire, 1961]. The
dotted lines in Figure 2.3 indicate the two regions in the concentration proﬁle of B-

kinetics based on diffusion in a bulk material.

 

1 For C-kinetic diffusion, only the material residing in the grain boundaries themselves contributes to the
total amount of diffusing species. In the case of diffusion barriers with low lattice diffusivities, grain-
boundary diffusion may be treated as the only diffusion mechanism if the grain boundary itself has a

thickness greater than 10 times of the extent of lattice diffusion of 2" D,t [Gupta, 1978].

In thin ﬁlms, the solutions for diffusant concentration are more complex because
of the additional boundary condition imposed by the ﬁlm surface1 [Gilmer, 1976 a&b].
The diffusant concentration raises signiﬁcantly near the free surface. Compared with the
diffusion in a bulk material, the diffusant proﬁle in a thin ﬁlm has higher concentrations
and lower gradient at the grain-boundary diffusion dominated region near the ﬁlm free
surface. Near the source/ ﬁlm interface, the diffusant concentration proﬁle is similar to
that for diffusion in the bulk model because of the dominant of the lattice diffusion. In
Figure 2.3, the solid line indicates the difference between the diffusion in a thin ﬁlm and
that in a bulk material. The y-axis of the proﬁle uses a normalized distance (y/Yo) away
from the diffusion source (Y0 is the ﬁlm thickness). 0n the source/ﬁlm interface, y/Y0=0,

whereas on the free ﬁlm surface, y/Y0=I.

In ﬁlms with small grain sizes, the ratio between the grain size and the distance of
diffusion inﬂuences the diffusant concentration proﬁle [Gupta, 1978]. When the ratio is
relatively large, smaller grain size allows the diffusant ﬁelds in the lateral direction to
overlap inside the grains more easily, making it closer to A-kinetics grain-boundary
diffusion [Gupta, 1978]. The A-type diffusion may be found in thin ﬁlm materials with a
small grain size even if the lattice diffusivity (D,) is small. For instance, in a material
with a grain size of 0.1 pm, a D, of 5x 10'15 cmzls and a diffusion duration of one hour

correspond to a ratio of 1.2, which is at the junction of the A- and B-kinetics.

2. Interaction between SiC and Nickel Aluminides

A sharp gradient in chemical potential at the interface between the reinforcement

and matrix in composites is the driving force for diffusion during processing or

 

1 To ﬁnd a mathematical solution for diffusant concentration proﬁle in thin ﬁlms, an imaginary source was
used in order to maintain the free-surface boundary condition for thin ﬁlms [Gilmer, 1976 a&b].

application of the composites at high temperatures [Chou, 1985; Reedy, 1989]. Inter-
diffusion and/or diffusion-controlled interfacial reactions often degrade the
reinforcement. The reaction products in the interfacial region are often brittle, and thus
detrimental to composite strength [Broutman, 1974; Schoutens, 1982]. SiC-reinforced
nickel aluminide matrix composites are typical examples of high temperature composites

[Pope, 1987; Yang, 1989a].

SiC reacts with y-Ni3Al -based alloys at temperatures above 1073 K [Nieh, 1989;
Yang, 1989b]. Graphite and a ternary compound, Ni 5.4.41 [Si2, are formed in the reaction
zone. A layer of NiAl is also formed near the matrix side [Chou, 1991b]. The reaction is
controlled by the diffusion of Ni. Figure 2.4 schematically presents the reaction zone
between a SiC and a ’Y-NI3AI plate after a 23-hour diffusion bonding treatment at 1273 K
(0.76 Tm of y-Ni3Al). The width of the reaction zone is found to be approximately 100 to
150 um under these conditions [Chou, 1991b].

SiC also reacts with B-NiAl-based intermetallic alloys, but at higher temperatures.
Above 1573 K, a Ni-Al-Si ternary compound, Ni 14AlgSi2, and graphite are formed [Chou,
1991a]. In Figure 2.5, the interface region from this reaction is schematically illustrated.
An annealing treatment at 1573 K (0.82 Tm of B-NiAl) for 20 hours produces a 10 um
thick reaction zone [Chou, 1991a]. This reaction is controlled by the decomposition rate
of SiC in the presence of free Ni. Compared with the SiC/y-Ni3Al system, the thermal

stability of the SiC/B-NiAl system is about one order of magnitude better.

The interface zones formed in the aforementioned reactions have inherently poor
mechanical properties [Nieh, 1989; Yang, 1989b]. Hence, there is a need for a diffusion
barrier between SiC reinforcement and the nickel aluminide matrices for the successful

development of these composites.

3. Diffusion Barriers

Principles of Diffusion Barrier

Diffusion barriers are thin-ﬁlm layers used to prevent two materials from coming
into direct contact, thus preventing inter-diffusion and interactions between the two
materials [0hring, 1991]. The selection of the barrier material is largely determined by
the characteristics of inertness and low diffusivities. Ideally, the barrier materials should
be chemically inert, and have negligible mutual solubility and diffusivity at the highest

temperature of application with respect to other materials in the system.

Other factors are important in choosing the barrier material as well. These factors
include ease of fabrication, similar coefﬁcients of thermal expansion between the barrier
and its substrate, stress-state in the barrier after fabrication, and compatibility with other
processing steps. Practically, some of the requirements for an ideal barrier are difﬁcult to
satisfy or even mutually exclusive. As a result, less ideal systems have to be adopted
instead [Kattelus, 1988]. Based on to the operational mechanisms, diffusion barriers are

divided into three categories: stuffed, passive, and sacriﬁcial barriers [Nicolet, 1978].

The imposition of a diffusion barrier upon a thermodynamically unstable system
slows down the equilibration process. Nevertheless, the instability in the system is never
actually removed. Enhanced reliability is achieved with a diffusion barrier at the cost of

increasing structural complexity and added processing expense [0hring, 1991].

Candidates for Diffusion Barriers

Many high-temperature alloy coatings (Ni-, Fe-, Co— based alloys with Cr, Ti, Al,

V additives) are used as barriers for prevention of diffusion bonding. The protection

mechanism is based on adherent impervious surface ﬁlms of A1203, SiO2, Cr02, or a

spinel-type, which grow on high temperature exposure to air [Grayson, 1983].

Low lattice diffusivity is crucial for high-temperature diffusion barriers because
lattice diffusion becomes dominant at high temperatures. Candidates for high-
temperature diffusion barriers are often stable compounds with limited deviations from
their stoichiometries. They include the transition metal carbides (HfC, TaC, TiC),

nitrides (AlN, HfN, TiN), borides (TiB2), and oxides (A1203, Y2O3) [Walters, 1988].

A1203 as a Diffusion Barrier in SiC/y-Ni3AI System

A1203 is a diffusion barrier candidate for many high-temperature systems. The
advantages of A1203 as a diffusion barrier includes a high melting point (2345 K),
moderate thermal shock resistance, and good chemical stability at high temperatures in a
wide variety of atmospheres. A1203 has a reasonably good and nearly constant strength

that can be maintained at temperatures up to 1373 K.

A1203 has a moderately high coefﬁcient of thermal expansion (CTE) of 8x 106/K,
as compared with other ceramics like SiC (4.7x 106/K). The CTE of A1203 ﬁts in
between SiC and N i3Al (12x106/K), which may limit the internal stress due to the
thermal cycling when an A1203 ﬁlm is coated on SiC reinforcements in SiC-reinforced

Ni3Al ~based matrix composites [Grayson, 1983; Nourbakhsh, 1989].

Furthermore, the diffusion coefﬁcients of Ni in oxides are low. For instance, Ni
has a diffusivity ranging from 10'13 to 10"2 mzls in amorphous Si02 over a temperature
range of 1273 to 1473 K [Goshtagore, 1969]. Nieh [1988] has shown the ability of A1203

ﬁlm to function as a diffusion barrier between planar SiC and a Ni3Al-based alloy.

10

In Nieh's experiment, the nickel aluminide was IC-50 (Ni - 23 at% Al - 0.5 at%
Hf — 0.2 at% B) which was pre—oxidized to form an A1203 surface ﬁlm. The pre-
oxidation was performed at 1273 K in air for 60 hours, after which a 0.13 um thick A1203
scale was formed on the surface of the Ni3Al plate. Subsequent diffusion bonding tests
were conducted using the pre-oxidized aluminide plate and a SiC disc. No bonding was

found even at 1373 K.

4. Comparison of Aluminum Oxide and Aluminum Nitride

Aluminum oxide is an effective barrier to Ni diffusion, and has demonstrated its
potential as a diffusion barrier between SiC and a Ni3Al-based alloy [Nieh, 1989]. The
mismatch of the coefﬁcients of thermal expansion (CTE) between A1203 and SiC,
however, may induce large thermal stress1 in an A1203 ﬁlm deposited on a SiC substrate
when subjected to a large temperature excursion. A1203 has a CTE with an average value
of 8x 106/K in a temperature range from 298 to 2123 K. The CTE is 4.7x10'6/K for SiC
from 293 to 1173 K [Shaffer, 1964]. Thus, the calculated result of the thermally induced
stresses in the A1203 ﬁlm could exceed 1 GPa when the temperature of an A1203 coated
SiC changes between room temperature and 1273 K, assuming that the ﬁlm is still intact

and behaves elastically.

 

1 The thermal stress is induced by the difference between the deposition temperature and the application
temperature because of the difference in thermal expansion coefﬁcients of a ﬁlm and its substrate. The
thermal stress in a ﬁlm is estimated in one-dimensional approximation (neglecting the Poisson effect) by

0' ﬁlm = Eﬁlm (aﬁlm - asub )(Tdeposition _ Tapplication )

where E ﬂ", is Young's modulus of the ﬁlm, aﬂm and and, are the average coefﬁcients of thermal
expansion for the ﬁlm and substrate respectively. Tawny“, is the substrate temperature during deposition
and Tappumm is the temperature at application. A positive value of ”ﬁlm corresponds to a tensile stress.

11

Aluminum nitride, on the other hand, possesses an average CT E of 4.8x 106/K
from 293 to 2073 K [Poluboyarinov, 1980; Shaffer, 1964]. The closeness of the CT Es of
aluminum nitride and SiC makes aluminum nitride an attractive candidate as a diffusion
barrier material on SiC particles. The thermally induced compressive stresses in an
aluminum nitride ﬁlm on a SiC substrate are calculated to be 31 MPa when temperature
reaches 1273 K from room temperature, and 38 MPa when reaching 1673 K, given an
elastic behavior of the ﬁlm. Figure 2.6 compares the CTEs among aluminum nitride,

A1203, and SiC, from room temperature to 873 K.

Aluminum nitride has attracted interest in recent years. This material has many
useful properties, such as high thermal conductivity, low thermal expansion, low density,
and high speciﬁc modulus [Sheppard, 1990]. Aluminum nitride is chemically stable at
temperatures up to 1873 K with both [3- and y- nickel aluminides and SiC [Lowell, 1990;
Zangvil, 1985]. Aluminum nitride has a relatively low strength and modulus at high
temperatures, which makes aluminum nitride an unattractive reinforcement candidate by

itself for high-temperature nickel aluminide matrix composites.

Aluminum nitride has a high resistance to thermal shock [Kingery, 1976], which
is a desirable feature for ceramics undergoing a large temperature excursion. The
resistance to thermal shock may be measured by the maximum-quench-temperature
diﬂerence that the material is able to withstand [Kingery, 1976]. Compared with 150 K
for A1203 [Kingery, 1976], the resistance to thermal shock of aluminum nitride is
approximate 550 K [Poluboyarinov, 1980]. The better thermal shock resistance of
aluminum nitride is attributed to a lower elastic modulus, a lower CT E, and a higher
thermal conductivity for aluminum nitride comparing with aluminum oxide. Figure 2.7
compares the thermal conductivities among aluminum nitride, SiC, and A1203, in a

temperature range from 297 to 573 K.

12

Aluminum nitride ﬁlms have been deposited by a variety of methods, from
physical vapor deposition (PVD) to chemical vapor deposition (CVD) [0hring, 1991].
Among these techniques, ion-beam assisted PVD methods offer attractive features,

especially for compound ﬁlm deposition, as will be discussed in the following section.

5. Vacuum Evaporation and Ion-Beam Assisted Deposition
PVD Techniques for Thin Films

During physical vapor deposition (PVD), physical mechanisms (evaporation or
collision impact) are used to bring materials into a gas phase from solid or molten sources
[0hring, 1991]. PVD is capable of depositing virtually any type of inorganic, as well as
some organic materials, over a wide range of deposition rates and temperatures. PVD is
a more controllable process than other thin ﬁlm deposition processes, such as chemical

vapor deposition (CVD) [Bunshah, 1984].

One PVD technique is sputter deposition. Sputtering is a vacuum process in
which atoms of a target material are ejected by the bombardment of ions and energetic
particles generated by an electrical discharge, and then condensed onto a substrate
[Holland, 1974]. Sputter deposition offers many advantages, including reproducibility,
continuous nature, and simplicity in operation. Sputtering itself is a complex process,

and has been extensively studied [Behrisch, 1981, 1983 and 1991].

Sputter deposition offers depositing atoms with kinetic energies in the range of 5
to 30 eV. Such energetic atoms improves the adhesion of deposited thin ﬁlm to its
substrate [Westwood, 1988]. The deposition process usually occurs in a relatively high

pressure range of 10'2 to 101 Pa, which causes a relatively high degree of impurity

13

incorporation [Mattox, 1989; Ohring, 1991]. For an alloy or compound target, the ﬁlm

concentration can easily match the composition of the target with a little variation.

Vacuum Evaporation

Another widely used PVD technique is vacuum evaporation. In the evaporation
process, the coating material is evaporated in a vacuum environment by heating. The
evaporation process can offer a high deposition rate for virtually all kinds of materials
[Glang, 1970; Ohring, 1991]. Among various heating sources for evaporation, the
electron-beam source provides the possibility to deposit high purity ﬁlms [Bunshah,

1984; Ohring, 1991].

Vacuum evaporation is a relatively "clean" deposition process because it is
performed in a vacuum environment below 10'4 or 103 Pa. Nevertheless, gas molecules
in the vacuum environment will impinge and be adsorbed on the growing ﬁlm surface.
Using the arrival rate of depositing atoms (10) on substrate surface and the gas
impingement rate (18), the concentration of the impurity molecules (C8 in at%) is

approximated as [Chambers, 1989; Ohring, 1991]:

J M P
g=——g——z4.38x10‘4-—a—
Ja+Jg sp M T

8
Here, 3 is the deposition rate in cm/sec, p is the ﬁlm density in g/cm3, P is the partial
pressure in Pa, Ma is the molar weight of the ﬁlm material in g/mole, M g is the molar

weight of the gas molecule in g/mole, and T is the background temperature in Kelvin.

Because the sticking probability of a gas molecule to a ﬁlm surface is usually less
than unity [Ohring, 1991], the above equation overestimates the amount of impurity

molecules residing in the ﬁlm by assuming the sticking probability to be unity. Using a

14

density value for a bulk material, on the other hand, the impurity concentration may be
underestimated by the above equation [Ohring, 1991]. For metallic materials, the ﬁlm
densities can be about 95% of their bulk value. Compound ﬁlms usually have lower

ratios, such as 70% for ﬂuorides [Pulker, 1984].

Conformal coverage (the coverage of the steps on substrate surface) is another
issue in vacuum evaporation. Evaporated atoms travel in essentially collisionless paths
prior to condensation on the substrate [Bunshah, 1984]. Because the evaporated particles
usually come from a point source, shadowing effect may cause unequal deposition
between the top and sides of the step on substrates. Inadequate step coverage, or
sometimes no step coverage at all, can lead to minute cracks in the ﬁlm [Ohring, 1991].
Such problems may be solved by either physically exposing the shadowed area to the
coating material ﬂux during deposition, or improving the depositing atoms' mobilities on
the substrate surface. The latter approach can be implemented by elevating the substrate

temperature [Thomton, 1977].

Evaporating compounds or alloy sources may form ﬁlms with compositions that
differ from that of the source. The constituents of alloy sources evaporate almost
independently of each other. Decomposition or dissociation occurs in compound
evaporation [Bunshah, 1984; Ohring, 1991]. Additional material sources may be needed
to form a compound or alloy ﬁlm with the same composition as the primary evaporating
source. For example, introducing oxygen ﬂow into the evaporation environment is a
common practice necessary to overcome oxygen deﬁciency in ﬁlms deposited by

evaporating a stoichiometric A1203 source.

Some compound ﬁlms can also be deposited by evaporating pure metal
evaporants in a relatively high partial pressure (0.1 - 5 Pa) of reactive gases, such as 02,

N2, CH4, C2H2, etc. [Ohring, 1991]. A plasma discharge may also be used within the

15

reactive zone between the metal source and substrate, helping to overcome the energy
barrier of reaction by ionizing both the metal vapor and the reactive gases. The plasma
discharge increases the reactivities of the coating particles on the surface of the growing
ﬁlms, and promotes stoichiometric compound formation [Bunshah, 1987; Ohring, 1991].
Compounds can be synthesized by this method include oxides (A1203, V203, TiO2),
carbides (TiC, ZrC, HfC, VC), and nitrides (TiN, MoN, HfN, BN).

Vacuum evaporation is a capable ﬁlm deposition process by itself. Nevertheless,
low kinetic energy associated with the evaporant atoms (about a fraction of an eV) results
in a weak adhesion of the condensed ﬁlms to their substrates. Also, decomposition or
dissociation that occurs during compound and alloy evaporation may make it difﬁcult to
form stoichiometric compounds or alloy ﬁlms. Ion-beam techniques are utilized to

improve the evaporated ﬁlm quality, as well as to form compound ﬁlms.

Reactive IBAD for AlN Films

Introduction To improve the quality of a deposited ﬁlm from evaporation
and to form compound ﬁlms, a hybrid technique has been developed which uses ion-
bearn technology with vacuum evaporation. It is called ion-beam assisted deposition

(evaporation), or IBAD.

The IBAD process provides independently controlled deposition parameters and
the characteristics of ion bombardment on substrates [Harper, 1982 and 1983; Ohring,
1991]. During the process, ion bombardment on a growing ﬁlm heightens diffusion and
chemical reactivity of ﬁlm atoms, and lowers the sticking probabilities of residual gas
atoms. With proper ion-beam parameters, ion bombardment during deposition can alter

coating stresses, and improve step coverage and packing density. In many cases, low

16

beam energies (about 100 eV or lower) and low ion-to-atom arrival ratios (a few percent
or lower) are most effective [Harper, 1983]. Figure 2.8 shows an IBAD conﬁguration

that incorporates an ion source and an electron-beam evaporation source.

The IBAD technique has many advantages similar to those of the sputtering
technique, because of the ballistic nature of the bombardment of energetic particles on
substrates and growing ﬁlms. In addition, the energetic bombarding particles can be
either ionized or neutral, which makes the IBAD technique applicable to non-conducting
substrates or ﬁlms. The technique is medium-to-high vacuum compatible also results in

less impurity incorporation into the growing ﬁlm.

The Ion Source In the heart of many IBAD processes is the broad-beam
(Kaufman) ion source [Harper, 1983; Kaufman, 1982; Ohring, 1991]. The Kaufman
source operates in a background pressure of approximate 10'2 Pa that is compatible to
both sputtering and evaporation. The resulting beams have a low-energy spread
(typically 10 eV) and are well collimated (with divergence angle of a few degrees). Such
ion sources have been widely employed in sputtering, etching, surface compound

formation, and thin ﬁlm deposition [Harper, 1982 and 1983].

The ion beam from a Kaufman-type source consists primarily of a mixture of
single- and double-charged molecules. Triple-charged (or even higher) gas ions are rare
under normal operational conditions of the source [Kaufman, 1987]. For nitrogen (N2)
gas, the single-charged species are N g and N + , and the double-charged are N2” and
N H . Among these four species, one N 5' + molecule and a pair of N + ions are

identical in terms of the number of atoms, momentum and energy.

17

Because the energy required to form a single-charged molecule is much lower
[Harper, 1982; Kaufman, 1987], the probability of forming a single-charged molecule is
much greater than the probability of a double-charged one. Consequently, the number of
double- charged species such as N H from a Kaufman source may be negligible. Taylor
[1981] has found that approximately 96% of the nitrogen ions are N 5' and 4% of N + .
Netterﬁeld [1988] has measured that the ratio of N; : N + emitted from a Kaufman ion
source is 3.7: 1 under a normal operational conditions at a 5x10‘3 Pa working pressure,

which indicates 78.7% of N3" and 21.3% of N+.

Reactive IBAD Using a reactive gas, such as N2 in the ion source, the
Kaufman source produces a reactive ion beam which can be used to form compound
ﬁlms with an evaporating metal [Harper, 1982]. In such a process, the degree of
incorporation of the reactive components in the growing ﬁlm is the key parameter that
controls stoichiometry, density, resistively, and many other properties of the ﬁlm
[Harper, 1983; Stelmack, 1989]. The atomic ratio between the arrival rate of the charged
species in the ion beam and that of the metal atoms is the parameter that, in practice,
affects the ﬁlm properties the most. Hentzell [1985] has shown that a ratio of N/Al 21
was necessary in order to form a stoichiometric aluminum nitride ﬁlm without metallic

aluminum grains, where the nitrogen atoms were assumed to be as N; .

Aluminum Nitride Film Deposition with Reactive IBAD In the reactive
IBAD process for aluminum nitride ﬁlm deposition, aluminum atoms are supplied by a
metallic aluminum source whereas the growing ﬁlm is bombarded with a nitrogen ion
beam [Harper, 1982; Hentzell, 1985; Netterﬁeld, 1988]. The charged nitrogen ions are

necessary for nitride formation because metallic aluminum atoms do not readily react to

18

neutral nitrogen molecules by chemisorption during depositions [Taylor, 1981]. The
reaction starts, however, when the nitrogen molecules become charged [Hentzell, 1985;
Netterﬁeld, 1988; Taylor, 1981]. The sticking coefﬁcient of nitrogen ions on aluminum

is then close to unity [Harper, 1985].

The atomic arrival ratio of nitrogen to aluminum (the N/Al ratio) is critical to the
properties of the deposited ﬁlm [Harper, 1985 ; Hentzell, 1985]. There are well-deﬁned
grains of metallic aluminum and aluminum nitride existing even when the N/Al ratio is
0.96. Metallic aluminum grains will only disappear after the N/Al ratio reaches 1.0 (In

their reports, it was assumed that the nitrogen ion beam consisted of 100% of N3" ions).

The microstructure of deposited aluminum nitride ﬁlms have a strong dependence
on the energy of the nitrogen ion beam [Hentzell, 1985]. At low beam energy (100 eV),
the ﬁlms have well-deﬁned ﬁber-type grains, whereas the grains are large and equiaxed
under a 500 eV nitrogen ion beam. The grain size increases with the ion beam energy.
The preferred grain orientation also depends on the beam energy. In low beam energy
regime (S 200 eV), the c-axis of aluminum nitride lattice is normal to the ﬁlm plane. The

c-axis becomes parallel to the ﬁlm plane when the beam energy is above 400 eV.

The visual appearances of aluminum nitride ﬁlms correlates with the ﬁlm
microstructures [Hentzell, 1985]. The ﬁlm appears metallically shiny when the N/Al
ratio is between zero and 0.54. When N/Al ratio is higher than 0.54, the ﬁlm becomes
dull gray until the ratio reaches 0.82. Between 0.82 and 1.0, the ﬁlm is shiny gray. A

pure aluminum nitride ﬁlm (N/Al = 1.0) is transparent.

The electrical resistivity is dependent on ﬁlm structure. The characteristic of the
resistivity changes from a metallic aluminum ﬁlm to a dielectric aluminum nitride ﬁlm
according to the N/Al ratio. In Hentzell’s experiment [1985], the resistivity of the ﬁlm

became too high to be easily measured with their equipment when N/Al exceeded 0.75.

19

Thin Film Deposition on Particulate Substrates

In PVD techniques, line-of-sight transport is followed in dealing with the
positioning of the substrate relative to the source. The location and orientation of a
substrate relative to the source not only determines the ﬁlm deposition rate, but also

strongly inﬂuences the ﬁlm structure and properties [Ohring, 1991].

Techniques have been developed for thin ﬁlm deposition on particulate substrates
[Caims, 1972]. It has shown that small particles can be levitated by vibrating their
container during deposition to achieve uniformly deposited thin ﬁlms on the particulate

substrates [Hillegas, 1990].

A1203 ﬁlms have been deposited on SiC particles with the help of a particle
levitation system [Cai, 1992]. Using an acoustic wave of 100 to 140 Hz in frequency,
SiC particles were levitated in their container during deposition. The ﬁlm deposited on
the particulate substrates in the particle levitation system had a fairly uniform thickness

and coverage.

6. Zone Structures of Thin Films from PVD

The ﬁlm structure is greatly affected by the homologues substrate temperature
(Ts) [Thomton, 1977]. The structures of a ﬁlm deposited by sputtering or evaporation
have been classiﬁed in terms of "zone structures". Three "zone structure" types have
been used for describing electron-beam evaporated ﬁlm morphologies [Movchan, 1969].
Each zone has distinguishable features, resulting from different ﬁlm nucleation and
growth mechanisms. Figure 2.9 shows schematically the zone-structure in evaporated

thin ﬁlms [after Movchan, 1969].

20

The structural differences among different zones are the result of differences in
atomic mobilities on the substrate surface [0hring, 1991]. At low substrate surface
temperatures (below 0.3Tm), the low atomic mobility both on the surface and inside ﬁlm
causes virtually all grain boundaries in the ﬁlm to be immobile. Hence the ﬁlm is formed
at this temperature regime by continued re-nucleation of grains during deposition. The
resultant ﬁlm consists of small tapered crystals with voided boundaries. Such a ﬁlm

structure is termed as the Zone 1 structure.

The Zone 11 ﬁlm structure is formed because the grain boundaries are mobile at
the substrate temperature range of 0.3 to 0.5Tm. Granular epitaxy and subsequent grain
growth result in a ﬁlm structure with columnar grains and dense grain boundaries. At
higher temperatures, enhanced surface diffusion results in Zone III structure with large

equiaxed grains.

A transition ﬁlm structure, Zone T, is formed at the high-end of the Zone I
temperatures because of the enhanced atomic mobility at this temperature regime. Films
deposited in this temperature regime are produced by continued nucleation of grains
during deposition and subsequent grain growth. Hence, the Zone T structure has a
mixture of small and large grains. Large residual stresses could be formed in the Zone T

region [Thomton, 1977 and 1989].

The grain sizes in evaporated ﬁlms vary with the structure of the ﬁlm. Hentzell
[1985], et. al., found that the grain sizes in 0.1 um thick evaporated metal ﬁlms ranged

from 0.01 um for Zone I structures to approximately 1 pm for the Zone 111 structures.

The temperatures dividing the structural regimes are affected by other factors as
well [Ohring, 1991; Thornton, 1977]. Roughness on substrate surface and an oblique
deposition angle promote the Zone 1 structure. An elevated inert gas pressure at low

temperature regime causes grain boundaries to become more open in both sputtered and

21

evaporated coatings. Such an inﬂuence largely vanishes at elevated temperatures. Ion
bombardment suppresses the Zone I structure, resulting in a structure similar to the Zone
T type. In an IBAD process, the ion-beam energy affects the structure of the ﬁlm.
Hentzell [1985] reported that the aluminum nitride grains became larger at higher ion
beam energy (500 eV). The grains tended to be equiaxed grains at high beam energy,

whereas elongated grains resulted at lower beam energy (100 eV).

7. Grain Growth and Solid-state Sintering in Ceramics

The structure and properties of deposited thin ﬁlms can evolve further under
application conditions. Such an evolution in ﬁlm structure and properties can have

signiﬁcant impact on the functionality of the ﬁlms.

Three major changes usually occur on heating of ﬁne-grained ceramic materials:
there is an increase in grain size; there is a change in pore shape; there is change in pore
size and number [Kingery, 1976]. At elevated temperatures, two transformation
processes in the microstructures of ﬁne-grained, porous ceramics usually take place,

namely solid-state sintering and grain growth.

Solid-state sintering usually occurs in porous ceramic materials. In a porous
ceramic material, both the pore size and shape change upon annealing. The pores
become more spherical in shape and smaller in size as the annealing continues. Changing
of free energy that gives rise to densiﬁcation of the ceramic is associated with a decrease
in surface area, and lowering of the surface free energy by elimination of solid-vapor
interfaces. Lowering of the free energy usually takes place with the coincidental
formation of new but lower-energy solid-solid interfaces [Kingery, 1976]. The rate of

sintering is roughly proportional to the inverse of the grain size.

22

The grain growth process is more often associated with ceramics with an
originally dense structure. In such materials, a reduction of free energy can be achieved
by increasing the average grain size. As the average grain size increases, some grains
shrink and disappear. The rate of growth increases exponentially with temperature, and
the activation energy corresponds approximately to the activation energy for grain-

boundary diffusion [Kingery, 1976; Smallman, 1985].

During grain growth in ceramics, pores on the grain boundaries may be left
behind by the moving boundary or migrate with the boundary, gradually agglomerating at
grain corners [Kingery, 1976]. In the early stages of annealing, a cluster of small pores in
the corner of a grain is a commonly observed result. In the later stages of grain growth,
when the grain size is larger and the driving force for boundary migration is low, it is

more usual for pores to be dragged along by the boundary, slowing grain growth.

23

CHAPTER 3. EXPERIMENTAL PROCEDURES

In this study, aluminum nitride and aluminum oxide ﬁlms were evaluated as
diffusion barriers in composite systems of SiC particle reinforced B-NiAl and y-Ni3Al
matrices. This study included the following tasks: to develop a heated particle levitation
system; to deposit ﬁlms on particulate substrates at various elevated temperatures; to
characterize the deposited ﬁlms; to study ﬁlms’ thermomechanical behaviors on SiC
particles; and to examine the effectiveness of both ﬁlms as diffusion barriers in the

composite systems.

1. Materials

In this study, model composites were used to study the effectiveness of the
diffusion barrier materials. Abrasive-grade SiC powders were chosen because of their
relatively large size. SiC particles of a nominal size of 60 um were purchased from
Mager Scientiﬁc Inc. Figure 3.1 shows a SEM photo of an as-received SiC particle. The

surface of the particle was characterized as ﬂat cleavage surfaces joined by sharp edges.

Both B-NiAl and y-Ni3A1 plates were used as composite matrices in this study.
They were at their stoichiometries (10.5 at%). The y-Ni3Al was doped in Oak Ridge
National Lab (ORNL) with 0.25 at% of boron to enhance its room temperature ductility.

The B-NiAl was donated by NASA.

Both aluminum nitride and aluminum oxide ﬁlms were deposited using ion-beam
assisted deposition (IBAD) process. A1203 rods of 6.3 mm in diameter (with 99.5%

purity) were used as the feedstock in a 3-kW electron beam evaporator to form the oxide

24

ﬁlms in an IBAD process. Argon gas of 99.99% purity was used to produce an assist
ion-beam in a 3-cm Kaufman ion beam source (the Kaufman gun). 99.99% pure oxygen
gas was leaked into the deposition environment as an oxygen content supplement for the
oxide ﬁlm depositions. For the nitride ﬁlm depositions, commercially pure aluminum
rods (99.0 wt% aluminum) of 6.3 mm in diameter were used. Nitrogen gas of 99.99%

purity was used to produce a reactive ion assist-beam from the Kaufman gun.

Besides the SiC particulate substrates, planar substrates were also used for ﬁlm
characterization. Glass microscope slides and polished Si wafers of 50.8 mm in diameter
(purchased from PureSil, Inc.) were used for ﬁlm composition and morphology studies.
Polished y-Ni3Al, B-NiAl and monolithic aluminum nitride plates were used for a cursory
study of Ni diffusion in the nitride material. The aluminum nitride plates were donated
by the Carborundum Company (Sanbom, NY) and had a nominal composition of 97.5%

aluminum nitride, and 2.5% Y2O3 in balance.

2. Particle Levitation and Heating System

System Overview

To facilitate the IBAD process for particulate substrates, a deposition system was
specially designed to achieve an uniformly deposited ﬁlm on the particulate substrates.
The system was designed based on the idea that particles can be levitated by vibrating
their container at certain frequencies [Caims, 1972; Hillegas, 1990]. In this study, a
particle levitation system was designed to accommodate ion-beam assisted depositions at
elevated temperatures. Radiant heaters were used to sustain the temperature on substrate

surface during depositions.

25

There were three major components in the particle levitation system, as
schematically shown in Figure 3.2. A particle holder assembly was located in the
deposition chamber. A vibration generator system was outside of the vacuum chamber,
which converted a sine wave signal from a signal generator into a mechanical vibration.
The third component of the levitation system was a connector assembly communicating
the holder and the vibrator. Proper cooling procedures and radiant shielding were
designed to prevent the system and the vacuum chamber from overheating at elevated

deposition temperatures.

In the next section, a detailed description of the three components of the levitation
system will be given. Later, a heater assembly will be described, which was added to

complete the system used in this study.

Levitation System

Particle Holder Assembly The holder assembly consisted of a quartz
container and two stainless steel brackets. The container was a 76 mm long, 51 mm wide
and 13 mm deep box with a maximum deposition area of 39 cm2. The two brackets,
which were made of 0.8 mm thick stainless steel, were locked onto four quartz stubs
attached to the four comers at the bottom of the container. The brackets were also locked
onto the quartz tube in the connector assembly (which will be described later in this
section). The sturdy and yet ﬂexible brackets provided a combination of friction ﬁt,

spring tension, and a good vibration transfer.

Vibration Generator System The central element of the generator system

was a 6-inch, 8 Q impedance audio loudspeaker (the acoustic transducer) which

26

generated a mechanical movement from an electronic signal. A quartz coupling was
designed to connect the speaker to the quartz tube in the connector assembly. The
coupling was attached to the speaker at speaker’s axis, using a resin epoxy. A thin
aluminum plate was used as a reinforcing rib between the coupling and the speaker to

enhance the vibration efﬁciency and reduce warping of the transducer.

A frequency adjustable electronic signal generator was used to obtain a sine wave
with an optimum frequency that could levitate particles with a limited bouncing effect.
The sine wave was ampliﬁed to an appropriate input signal power level to the speaker.
The optimum frequency was inﬂuenced by the characteristics of the levitation system and
the particles involved. It was also a result of vacuum and deposition conditions. For the
conditions used in this study, the optimum frequency for a sine wave signal was found in
a range of 100 to 140 Hz. A 2 to 3 Watt power input was necessary for the particle

levitation system to operate with a load of one gram of SiC particles.

Connector Assembly The connector assembly served as a communicator
between the particle holder in a vacuum chamber and the vibration generator system
outside. The assembly consisted of a quartz tube sealed at the inner end, an O-ring
vacuum chamber sealing component, and a weight balancing mechanism. A coaxial air
tube, 6.4 mm in diameter, was used for forced cooling of the interior of the quartz tube

for elevated temperature applications.

The quartz tube of 38 mm in outer diameter was inserted through a 38.1 mm
vacuum chamber port assembly with a Viton O-ring seal. As the fulcrum of the whole
vibration system, a pin in the tube was locked on the cover of the port assembly. The
fulcrum pin secured the orientation of the particle holder during depositions. The vertical

movement of the tube at the O-ring position was negligibly small because of the short

27

distance (about 1 mm) from the O-ring to the fulcrum pin. There was practically no air-

leakage found during the depositions in this study.

To achieve a balanced levitation system around the fulcrum pin, spring tensions
were applied to the quartz tube near its outside end. Two springs were used. Because of

the balance, there was little strain on the vibration generator.

Heating Zone

To maintain an elevated temperature on substrate surfaces during deposition, two
heater compartments were placed symmetrically near the specimen holder. Figure 3.3
illustrates the construction of the heater compartments. Each compartment consisted of a
heater enclosure and a water-cooled shield. The heater enclosure was made of 1.6 mm
thick stainless steel, capable of containing up to three quartz-halogen radiant heaters
(Sylvania BVE-type reﬂector heaters, rated 625 Watts at 110 Volts). An adjustable
power transformer supplied the power to the heaters. Using three heaters in each of the
heater compartments, the temperature in the particle holder could reach up to 973 K. The

water-cooled shield was made of 1.6 mm thick copper sheet.

The heater enclosures were mounted to the water-cooled shield with adjustable
set-screws, as shown in Figure 3.3. These screws maintained the height of the heaters
relative to the substrates as well as the angle of the radiation. The heaters were placed as

close to the substrates as possible to ensure heating efﬁciency.

28

Performance of the System

The levitation system functioned well at the frequencies normally used for
levitation (100 - 140 Hz). For one gram of SiC particles (with an average size of 60 pm)
in the particle container, a sine wave of 2 to 3 Watts to the vibration generator was
sufﬁcient for levitating the powders at an optimum frequency of 120 Hz in both ambient
and vacuum conditions. During deposition, this optimum frequency varied in a range of
100 to 140 Hz. During levitation, particles were observed to gather at locations with

minimum vibration magnitude inside the container.

The system was capable of sustaining at working temperatures up to 973 K in the
deposition area of 76 mmx51 mm. A deposition temperature of 793 K was achieved
using a total of 775 Watts input power. The temperature in the deposition area was
relatively uniform. At an average temperature of 793 K, the temperature variation in the

deposition area was measured at :20 K.

The performance of this system was stable over a long working period of tens of
hours. The system had been used more than 301evitation hours at a deposition
temperature of 793 K. The system was able to withstand thermal cycles between room

temperature and the deposition temperature without apparent damage.

This system has been used for depositing aluminum nitride and oxide ﬁlms at
various temperatures and on various substrates. Both inert and reactive IBAD processes
were compatible to the levitation system. On particulate substrates, the ﬁlm thickness

was found to be fairly uniform.

29

3. Film Deposition, System and Procedures

IBAD System

The conﬁguration of an ion-beam assisted deposition system, the IBAD system, is
schematically shown in Figure 3.2. An electron-beam evaporation source and a Kaufman
broad-beam ion source were arranged in a vacuum chamber, together with a particle
levitation system. During depositions, the down-ﬁring, rod-fed electron-beam
evaporation source supplied the coating material ﬂux, while the 3-cm Kaufman ion

source directed an assist ion-beam towards the substrate surfaces.

By using a cryo-pump, a base pressure less than 5x10’5 Pa was achieved in the
deposition chamber at room temperature. Maintained by radiant heaters, the temperature
of the substrate area during deposition was measured by a thermocouple placed near the
substrates. Because the thermocouple was placed at the same area as the substrates, the

heat ﬂuxes for both the thermocouple and the substrates were the same during deposition.

The 3-cm ion source was equipped with a divergent accelerator grid that produced
an ion beam with a divergence angle of 10 degrees. With a working distance of 10 cm
from the substrates, the effective ion beam coverage area was calculated of 33.4 cm2 at
the substrate location. The current of the incident ion beam was measured with a Faraday

probe during deposition.

In each deposition-run, one gram of SiC particles was used. During deposition,
the particle charge was observed to cover roughly 40% of the area in the particle
container of 39 cm2. Hence, the average size of the particle covered area, which equals
the total exposure surface area of the particulate substrates to the incoming depositing

ﬂuxes, was estimated to be roughly 16 cm2.

30

Aluminum Nitride Film Depositions

A reactive IBAD process was applied for alurrrinum nitride depositions. During
the deposition, metallic aluminum was evaporated from an electron—beam evaporation
source. Nitrogen ions with 100 eV beam energy were supplied by a 3-cm Kaufman ion

source. The nitrogen ion beam was also served as a ballistic assist-beam.

According to researches done by Kaufman [1987], Netterﬁeld [1988] and Taylor
[1981], N '5 ions constituted most of the ions generated in the ion source (from 78.7% to
96%). Other species in the ion beam were mostly N + ions. Only a negligible amount of
the ion species was multi-charged ions. To calculate the atomic arrival rate, the nitrogen
ion-beam was assumed consisting of only N ‘5 and N + molecules, with x% N g and y%

N + . The ion density (i) for an ion beam with a current density of I (in A/cmz) was

i (ions/cmzlsec)

_ 1.6x10—l9

and the arrival rate of nitrogen atoms at the substrate surface ( J N ) was consequently

1,, = (2 x x% + y%) = (2 x x% + y%)-i (atoms/cmZ/sec).

1.6x10"9

To achieve an atomic ratio of MA] of unity (i.e. JN = 1,31), the ratio between the

ion density (1) and the arrival rate of aluminum atoms (1,“), the i/a ratio, was thus

1
2x% + y%)

 

i/a = 1713,: i/JN=(

Based on the reference data [Kaufman, 1987; Netterﬁeld, 1988; Taylor, 1981], the

required i/a ratio was between 0.51 and 0.56.

31

The ﬁlm deposition rates on planar substrates were calculated by using step-test
samples on planar substrates. The step-test samples were produced by depositing ﬁlms
on partially covered glass slides. Using a Dektak II proﬁlometer, the ﬁlm thickness was

determined by measuring the step depth at the ﬁlm edges on the step-test samples.

The deposition conditions for planar substrates are tabulated in Table 3.1.
Conditions for the deposition of aluminum nitride ﬁlms on both planar and particulate
substrates were similar, except for the duration. For depositions at 593 K, an average i/a
ratio of 0.56 was detected with an i/a range between 0.52 and 0.60. On planar substrates,
the averaged ﬁlm deposition rate at 593 K was measured of 0.18 nm/sec, and the average
ion density of the ion beam was 4.9x 10l4 (ions-cm'z-sec'l). For depositions at 793 K, the
ion density was averaged of 3.4x10l4 (ions-cm'z-sec'l). An average ﬁlm deposition rate
of 0.14 nm/sec on planar substrates resulted in an average i/a ratio of 0.53. The range of

the i/a ratio was between 0.42 to 0.64.

Table 3.1. Deposition Rates and i/a Ratios of AlN Films Deposited on

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Planar Substrates.
Sample T (K) 10"“) t (min.) s E Ion Arrival Rate i/a ['75
1-1 0.16 15 0.18 4.5x1014 0.52
593 0.18 0.56
1-2 0.16 15 0.18 5.2x1014 0.60
2-1 ‘ 0.10 10 0.17 3.4x1014 0.42
793 0.14 0.53
2-2 0.20 30 0.1 1 3.4x 1014 0.64
T: Deposition Temperature. 1: Film thickness. t: Deposition Duration.
i/a : Ion-to-Atom Arrival Ratio. i/a : Average Ion-to—Atom Arrival Ratio.

2

Ion Arrival Rate: in (ion -cm- -sec-l ). 5‘ : Deposition Rate on Planar Substrate, in (um/see).

E : Average Deposition Rate on Planar Substrate, in (hm/sec).

32

Prior to deposition, the majority of the residual gas species in the chamber were
nitrogen, oxygen, and water vapor. The partial pressures of oxygen and H20 were both
about 1.3x10’5 Pa. The working pressures were 3x 10'3 Pa for depositions at 593 K, and
7x103 Pa for 793 K deposition.

Aluminum Oxide Film Depositions

A1203 ﬁlms were deposited at room temperature using an IBAD process with an
inert argon ion assist beam]. An A1203 rod was thermally evaporated. The argon ion
beam of 100 eV was used as a ballistic assist-beam. The working pressure in the
deposition chamber was 4x10“3 Pa. An oxygen ﬂow was leaked into the deposition
chamber to compensate for an oxygen deﬁcit normally associated with evaporation of
A1203 [Bunshah, 1984; Ohring, 1991]. The oxygen partial pressure before deposition
was 3x104 Pa due to this oxygen supplemental ﬂow. The deposition rate for planar
substrates was measured to be 0.46 nm/sec. The i/a ratio was maintained at 0.02, with an

ion density of 2.3x 10l3 (ions-cm'Z-sec").

Deposition Conditions on Particulate Substrates

In this study, the targeted ﬁlm thickness on SiC particles was 0.5 pm for all ﬁlms.
Because of the high surface-to-volume ratio of the particulate substrates, the effective
ﬁlm deposition rate on SiC particles was much lower than the rate on planar substrates
[Cai, 1992]. To achieve the targeted ﬁlm thickness, a total of 9.75 hours of deposition

duration was required for aluminum nitride ﬁlms deposited on SiC particles at 593 K.

 

l A reactive IBAD process for A1203 ﬁlm deposition was not employed because the Kaufman ion source
was not compatible with a high oxygen ﬂow required for such a process.

33

The duration was 25 hours for depositions at 793 K. For A1203 ﬁlm deposition, SiC

particles were coated for 6 hours to achieve the targeted ﬁlm thickness.

The deposition temperature was measured by a thermocouple placed at the
deposition area. The thermocouple was located so that it was subjected to the same
radiant heat and deposition ﬂuxes as the substrates. Thus, the temperature reading
reﬂected the temperature of the substrate surfaces that were exposed to both radiant heat
and deposition ﬂuxes. In this experiment, the reading of the thermocouple was assumed

to be the deposition temperature of the substrate surface.

For depositions on particulate substrates, the deposition area was static. However,
the substrates were levitated, moving in and out the deposition area dynamically. This
levitation of the substrates caused the cycling of the depositing surface on each
particulate substrate. This cycling not only interrupted the growth of the depositing ﬁlm,

but also cycled the temperature on the substrate surfaces.

With depositing ﬂuxes and a direct radiant heat ﬂux, the substrate surface
experienced the full deposition temperature and deposition rate. The deposition rate on
the surface was the same as the rate on planar substrates under the same deposition
conditions. When the surface was orientated away from the ﬂuxes, The ﬁlm deposition
stopped. The substrate surface gradually cooled down towards the surrounding
temperature. In this case, the surrounding temperature was presumably the temperature
of the quartz particle holder. The holder was measured to be approximately 423 K during
depositions. Because of the limited size of the deposition area (about 16 cm2) and the
large surface area of the particle charge (estimated of over 300 cmz), the particles spent
most of their time (about 95% of the total deposition time) in the non-coating stage

during the deposition period.

34

4. Film Characterization

The microstructure, thickness, and composition of the deposited thin ﬁlms were
studied in detail because of their importance in inﬂuencing the performance of the thin

ﬁlms as diffusion barriers.

Film Structure and Thickness

After deposition, the surface morphology of ﬁlms on both planar and particulate
substrates was examined in scanning electron microscopy (SEM, Hitachi S-2500C). The
SEM samples were coated with a thin layer of gold to reduce surface charging associated
with dielectric materials. Secondary electron images (with electron beam accelerate
voltages between 10 to 20 kV) were used of ﬁlm structure and thickness study. The

internal scale provided by the SEM was used for size measurement.

To measure the ﬁlm thickness on SiC particles, as-coated SiC particles were
embedded in resin and polished with diamond compounds to reveal the cross-section of
the ﬁlms. The polished samples were then gold coated and studied in SEM. The
magniﬁcation rates of the SEM were in a range between 25000 and 35000. Multiple
thickness measurements (more than 20 in each case) were made, and the results were

averaged to determine the ﬁlm thickness on the particulate substrates.

The measured ﬁlm thickness on particles from SEM images varies depending on
how the ﬁlm is fractured after sectioning. A perfectly sectioned ﬁlm along the polishing
direction (Case 1) produces the highest ﬁlm thickness measurement on SEM images. On
the other hand, the lowest ﬁlm thickness reading on SEM images arises when the ﬁlm is
fractured perpendicular to ﬁlm surface (Case 2). The following discussions focus on the

average value of the ﬁlm thickness from SEM images for both cases.

35

The general assumptions of the discussions are as follow. Particles are spherical

and of a radius of to. The ﬁlm thickness is uniform on the particles and of a value of d.

d<<r0_ SEM images are taken from a direction perpendicular to the polishing direction.

The last assumption is that there is no error in the measurement on SEM images.

 

 

SEM Detector

 

 

t, (appeared on SEM image)

3 i Polishing Line

  
   

Particle

""55, - ”...-"...
MFLLHL

Case 1. Film Sectioned along
the Polished Line.

 

SEM Detector

 

 

 

t2 (appeared on SEM image)

ﬂ Polishing Line
__ __ -1---

  

“Inna—cu“

Case 2. Film Fractured Normal
to Film Surface.

In Case 1, the measured ﬁlm thickness (t,) is a

function of the location of the polishing line (x):

tl (x) = J00 +d)2 —x2 -‘/r02 —x2

The average value of the thickness measurement is

 

averaged over an x range from 0 to r0:

 

 

 

It,(x)dx 2

Z= ° ... = Md) sin"(—'° )

Idx 2r0 ro+d

0

+ r0 d cos(sin"( r0 ))—M
r0 +d 4
d << r0
_ 2
tl z Mn—in zzd = 1.57d.
4r0 2

In Case 2, the measured ﬁlm thickness (:2) is a

function of the location of the polishing line (x):

 

d
r0+d

 

t2(x)= -‘/(r()+d)2-x2

The average value of the thickness measurement is

averaged over an x range from 0 to r0:

36

r

_ (Itz (x)dx
0

d(ro+d) . .1

r0

 

 

t2 = r = srn (_r0__) +icos(sin "( ))
° 2r0 r0 + 2 r0 + d
[dx
0
(1 << r0
.7, z My: z 5d = 0.79d.
4n, 4

Therefore, the ratio between the measured ﬁlm thickness from SEM images (I)
and the real ﬁlm thickness (d) on particles is estimated to be between 0.79 and 1.57,

depending on the way the ﬁlm fractured after polishing.

Composition Study

The composition of the oxide ﬁlm was studied with a Link AN 1000 Energy
Dispersive X-ray spectrometer (EDX) installed on a Hitachi S-2500C SEM. The EDX
equipped with a windowless detector. By using a bulk A1203 material as a standard, the
composition of a deposited oxide ﬁlm was measured with quantitative corrections

computed using the manufacturer’s ZAP—4 code.

A Perkin-Elmer’s PHI 5400 X-ray photoelectron spectroscope (XPS) was used for
composition studies of the ﬁlms used in this study. With X-rays exciting photoelectrons
from the surface of a sample, the XPS was a surface analysis technique that provides both
elemental and chemical bonding information [Moulder, 1992; Walls, 1898]. Appendix I

of this paper states the principles of the XPS technique.

The XPS unit used in this study was equipped with an argon-ion sputter gun for

surface contamination removal and ﬁlm composition proﬁling. The sputter gun was

37

mounted off the axis of the photoelectron collector. The sputtering rate with this ion gun

was measured of 0.3 nm per rrrinute for the aluminum nitride ﬁlm.

Because the shadowing effect among particles interfered with the removal of
surface contamination (mostly C0, C02 due to the exposure of air after ﬁlm deposition),

XPS study of the nitride ﬁlms was performed on planar substrates.

Because a range of ion-to-atom ratio (the i/a ratio, ranging from 0.42 to 0.64) had
been detected during the depositions of the nitride ﬁlms on the SiC particles, planar
samples for the composition study were made under the similar conditions. The ﬁrst
sample deposited at 793 K had an i/a ratio of 0.42. The i/a ratio for the second sample,
which deposited at 593 K, was measured at 0.60.

5. Thin Film Thermomechanical Behavior: Annealing Tests

Experiments were conducted to study the structural thermomechanical behavior
of both aluminum nitride and oxide ﬁlms deposited on SiC particles after heat treatments
in both ambient and vacuum conditions. SEM was used to search for evidence of any
cracks or voids formed in the ﬁlms after annealing. The SEM samples were gold-coated
for improving surface conductivity. Carbon-tape was used as the carrier for the particles

in SEM sample preparation.

Annealing at 1273 K in Air

The annealing was performed in air to study the capability of the aluminum

nitride and oxide ﬁlms to withstand a moderate annealing temperature of 1273 K. SiC

38

particles with aluminum nitride and oxide coatings were heated in air from room

temperature to the annealing temperature (1273 K), and subsequently air-cooled.

In addition, some SiC particles with aluminum nitride ﬁlm (coated at 593 K) were
quenched in liquid nitrogen (77 K) from the annealing temperature of 1273 K. This test
provided a greater temperature difference comparing with the air-cooling test. The

greater temperature difference induced a higher thermal stress in the ﬁlm.

Annealing of Aluminum Nitride Films in Vacuum

For annealing temperatures higher than 1273 K, the annealing tests were
performed under vacuum conditions. Aluminum nitride ﬁlms, deposited on SiC particles
at various temperatures, were tested under conditions similar to those expected for high-
temperature applications. The results from this experiment provided information about

the ﬁlms’ behaviors during composite consolidation processes.

The coated SiC particles were heated to the annealing temperatures (1573 K and
1673 K) in one hour in a Centorr Vacuum furnace at a pressure of 5x 10'3 Pa, and then
annealed for 4 hours. After cooling down to 673 K in 30 minutes, the samples were
subsequently cooled in 8 hours to reach room temperature in vacuum. The annealed

samples were examined with SEM.

SiC particles with a nitride ﬁlm (coated at 793 K) were also annealed at 1373 K

for one hour in vacuum. Samples were examined with SEM after the annealing.

39

6. Ni Diffusion in Aluminum Nitride, a Cursory Study

Because of limited information on Ni diffusing in an aluminum nitride matrix,
experiments were designed to examine the extent of Ni diffusing in the nitride at elevated
temperatures applied in this study. Conclusion drawn from these experiments was

informative to determine the barrier ﬁlm thickness used in this study.

Diffusion in Bulk Aluminum Nitride

Monolithic aluminum nitride and nickel alurrrinides (B-NiAl and y- Ni3Al) plates
were hot-pressed (diffusion bonding) at a vacuum level was of 5x10’3 Pa. A bonding

pressure of 10 MPa was applied, which ensured sufﬁcient contact between the plates.

To provide a better contact between the nitride and aluminide plates, the
contacting surfaces of aluminum nitride, B-NiAl and y-Ni3Al plates were polished with
diamond compounds prior to the hot pressing. With a y-Ni3Al plate, the aluminum
nitride plate was hot-pressed for 32 hours at 1473 K (0.88 Tm of y-Ni3Al). A diffusion
couple of an aluminum nitride plate and a B-NiAl plate was hot-pressed at 1673 K (0.88
Tm of B-NiAl) for 12 hours in vacuum. After the bonding procedures, the diffusion

couples were cut to reveal the interface between the plates.

Energy dispersive X-ray spectroscopy (EDX) was applied to examine the distance
of Ni diffusing into alurrrinum nitride lattice in the diffusion-bonded samples. Line-scan
proﬁles of the Ni element across the interface were used to mapping the extent of Ni
diffusing in aluminum nitride from the aluminide matrices. Because a typical EDX probe
had a spatial resolution about 1 um [Loretto, 1984; Russ, 1984], this spatial resolution

limited the lowest diffusivity that could be estimated under the proposed conditions.

40

For instance, with the conditions of the nitride/y-Ni3Al diffusion couple (32 hours
at 1473 K), the spacial resolution of 1 um determined a minimum diffusivity of
2
x2 (l x 10.4)

__ = = 2.2x10'l4 cmz/sec.
4t 4 x 32 x 3600

 

Dmin =

Here the diffusion distance (x) is assumed to be the EDX's spatial resolution of l um.

For the conditions of the nitride/B-NiAl diffusion couple (12 hours at 1673 K), the
detectable limit of the diffusivity of Ni was 5.8x 10'” cmzlsec. For difﬁsivities less than
the minimum values, the EDX method was not capable of measuring the diffusion
distance accurately. Thus, when the diffusion distance was less than 1 pm, only an upper

bound of diffusivity could be determined from the EDX method.

Ni Penetrating Aluminum Nitride Films

The penetration depth of Ni diffusing in an aluminum nitride ﬁlm was studied.
The nitride ﬁlm of 225 nm thick was deposited on a y-Ni3Al plate at the same conditions
of nitride ﬁlm deposited on planar substrates at 793 K, as described in Section 3. The
coated y-Ni3Al plate was annealed at 1373 K for 30 minutes in a vacuum of 5x 10“1 Pa.

The Ni concentration in the nitride ﬁlm was then studied with XPS proﬁling.

7. Effectiveness of Diffusion Barriers

The effectiveness of the diffusion barrier ﬁlms was examined under the different
chemical potential levels. Composites of SiC particle reinforced B-NiAl and y—Ni3Al

matrices were selected for the low and high chemical potential conditions, respectively.

41

The survivability of the SiC particles in the composites during the compositing processes

was used as the gauge for evaluating the effectiveness of the diffusion barrier ﬁlms.

Coated and non-coated SiC particles were sandwiched and hot-pressed uniaxially
between two aluminide plates in a vacuum of 5x 10'3 Pa. A hi gh-temperature mechanical
testing system (MTS 810) with a Centorr vacuum furnace was used for the hot—pressing

process. The conﬁguration of the hot-pressing process is shown in Figure 3.4.

The hot-pressing conditions for the B-NiAl matrix composite were 4 hours at
1673 K under 10 MPa compressive stress. Conditions of one hour at 1373 K under 40
MPa compressive stress were used for the y-Ni3Al matrix composites. Before the hot
pressing, specimens were ramped in an hour to the process temperatures from room
temperature. Samples were cooled in vacuum after the hot pressing for eight hours to

reach room temperature.

SEM/EDX was used to study the interface of SiC and the aluminide matrices.
The B-NiAl matrix composite samples were fractured compressively at room temperature
to reveal the SiC particles. A sample of SiC particles with nitride ﬁlms coated at 593 K
were also electropolished to study the barrier ﬁlms’ conditions after the consolidation

process. The polishing agent was a solution of 10% perchloric acid and 90% acetic acid.

The y-Ni3Al matrix composite samples were polished and etched to reveal the
interface region. The etchant solution consisted of 5 grams of CuSO4, 20 ml of HCl, and
20 ml of H20. The etched samples were studied with an optical microscope. The sample
with SiC particles with the nitride ﬁlm coated at 793 K was an exemption, which was

fractured in tension at room temperature. This fractured sample was studied using SEM.

The next chapter presents the results from all the experiments described above.

Discussions will be given for the results. Conclusions drawn for this study will follow

42

CHAPTER 4. RESULTS AND DISCUSSION

1. Film Characterization

Three characteristics of thin ﬁlms were studied, which were ﬁlm thickness,
structure, and composition. These characteristics strongly inﬂuenced the effectiveness of
the ﬁlms as diffusion barriers. Furthermore, these features characterized the ﬁlm

deposition processes used in this study.

Film Thickness and Deposition Rate on SiC Particles

By polishing the coated particles that were embedded in resin, cross-sections of
the ﬁlms were exposed. After gold-coated, the polished sample was studied in a SEM.
Film thickness was estimated with an internal scale provided by the SEM imaging.
Figure 4.1 shows a SEM image example that captured the cross-section of an aluminum
nitride ﬁlm on a polished SiC particle. Measured directly on Figure 4.1 (b), a SEM

image with a 25000x magniﬁcation, the averaged ﬁlm thickness was 0.51 pm.

The same procedure was used to estimate the average ﬁlm thickness of the
alurrrinum oxide, aluminum nitride ﬁlms on their SiC particulate substrates after
depositions. It was estimated that the oxide ﬁlms had an average thickness of 0.47 um
after a 6-hour deposition, corresponding to an effective deposition rate of 0.022 nm/sec.
The nitride ﬁlms deposited for 9.75 hours at 593 K were of 0.52 um thick. The nitride
ﬁlm coated at 793 K was 0.54 pm in thickness after a 25-hour deposition. The effective

deposition rates for the nitride ﬁlm deposited on SiC particles at 593 K were thus 0.015

43

nm/sec, and 0.006 nm/sec for the nitride ﬁlm deposited at 793 K. Table 4.1 lists the

results of the ﬁlm thickness on SiC particles.

The actual ﬁlm thickness on particles could be different from the value measured
from SEM images. As discussed in Chapter 3, the magnitude of the difference was
determined by the way the ﬁlm was fractured after polishing. The ratio between the
measured ﬁlm thickness from SEM images and the actual ﬁlm thickness on particles was
estimated between 0.79 (when ﬁlm fractured normal to the ﬁlm surface) and 1.57 (when

ﬁlm fractured along the polishing line). The average value of these two ratios was 1.18.

In this study, ﬁlms were observed to have a tendency of fracturing normal to the
substrate surfaces. Hence, the ratio between the measured and the actual ﬁlm thickness
on particles was estimated between 0.79 and 1.18. By this estimation, the uncertainty
was about :l:19% when using the measured value as the actual ﬁlm thickness. Therefore,
after considering the errors and the uncertainty of the measurement technique, the

average values of ﬁlm thickness of all three ﬁlms deposited on SiC particulate substrates

Table 4.1. Measured Thickness of As-Deposited Films on SiC Particles.

 

Material T(K) t(hour) Anghk MaxThk MinThk §(nmlsec) 3. (nm/sec)

 

 

 

A1203 300 6 0.47 mm 0.6 pm 03 “m 0.022 0.024
AlN 593 9.75 0.52 um 07 um 03 um 0.015 0.009
AlN 793 25 0.54 pm 07 pm 04 um 0.006 0.007

 

 

 

 

 

 

 

 

 

 

T: Deposition Temperature, in Kelvin. t: Deposition Period. Anghk: Average Thickness.
MaxThk: Maximum Thickness Measured. MinThk: Minimum Thickness Measured.

3': Average Measured Deposition Rate on Particulate Substrate.
5‘8 : Estimated Deposition Rate on Particulate Substrate (= 0.052 5‘ pl ).

44

were 0.5 mm, with an error bar of i0.1 um.

With similar deposition conditions, as stated in Chapter 3, the average deposition
rates on planar substrate were 0.46 nm/sec for the oxide ﬁlm, 0.18 nm/sec for the nitride
ﬁlm deposited at 593 K, and 0.14 nm/sec for the nitride ﬁlm deposited at 793 K,
respectively. For an estimated ratio of 0.0521 between the effective deposition rate on
particulate substrate (Se) and the rate for the planar substrate (S p, ), the estimated
effective deposition rate for the oxide ﬁlm on SiC particles was calculated to be 0.024
nm/sec. The estimated effective rate was 0.009 nm/sec for the nitride ﬁlm deposited on
SiC particles at 593 K, and 0.007 nm/sec for the ﬁlm deposited on SiC particles at 793 K,

respectively. Table 4.1 summarizes the results of the ﬁlm thickness on SiC particles.

For the nitride ﬁlms deposited at 793 K and the A1203 ﬁlms, the measured
effective deposition rates on SiC particles were comparable to their estimated values.
However, the measured rate of the nitride ﬁlm deposited on SiC particles at 593 K was

0.015 nm/sec that was 167% of the estimated value of 0.009 nm/sec.

Because the effective deposition rates for the nitride ﬁlm deposited at 593 K was
much higher than the estimated rate, the ﬁlm on SiC particles was probably not as dense
as the ﬁlm deposited on the planar substrate. A different ﬁlm structure was likely to be
present on the particulate substrates compared with that on the planar substrate. Studies

on ﬁlm structures, discussed next, support this conclusion.

 

l The estimated ratio (r) between S, and Sp, is calculated as r = (Adp)/( 6 W), where W is the amount of

the particle charge used in a deposition, d is the nominal size of the particles, p is the density of the
particles, and A is the deposition exposure area [Cai, 1992]. In this study, W was 1.0 gram, (1 was 60 um,
A was in an average of 16 cm2, and p was 3.22 g/cm3. Hence, this ratio was estimated to be 0.052.

45

Film Structure and Morphology

Aluminum Nitride Film Deposited at 593 K Aluminum nitride ﬁlms
deposited on planar substrates at 593 K had dense structures with smooth, reﬂective
surfaces. At 0.4 um thick, the ﬁlms were transparent and non-conductive. Figure 4.2 (a)
shows the surface of the ﬁlm deposited on a glass slide. Figure 4.2 (b) shows the ﬁlm’s
cross-section. The ﬁlm consisted of small crystallites with a size of 0.1 — 0.2 pm. The

structure was classiﬁed as the Zone T structure.

0n the SiC particulate substrates, however, the nitride ﬁlm had a different
structure. After a 9.75-hour deposition, the nitride ﬁlm was formed with a globular
structure of grain clusters on the SiC substrate surface. Figure 4.3 shows SEM images of

ﬁlm morphology at both low and high magniﬁcations.

The globular clusters in the ﬁlms on SiC particles were of 0.3 pm in average size,
with visible voids in between the clusters. The globular feature of the ﬁlms resembled
the Zone I structure [Movchan, 1969]. Within each globular cluster, grains were densely
packed. The grains were of different sizes (0.2 pm or smaller). A Zone T structure was a
good description for the structure within each globular cluster [Movchan, 1969].
Therefore, the overall ﬁlm structure was determined as a mixture of Zone I and Zone T

SII'UCIUI'CS .

The large number of voids associated with the globular structure may explain the
difference between the measured effective deposition rate on particulate substrates and
the estimated one. It may be estimated, by using the difference in the estimated and the
effective deposition rates, that the void space in the as-deposited ﬁlm was about 40% of

the total ﬁlm volume.

46

Aluminum Nitride Film Deposited at 793 K Unlike the nitride ﬁlms on SiC
particles deposited at 593 K with a globular structure, the nitride ﬁlms deposited at 793 K
had a dense and smooth morphology on both planar and particulate substrates. Figure 4.4
shows SEM images of the ﬁlm morphology on SiC particles after a 25-hour coating. A
cross-section of the ﬁlm was shown earlier in Figure 4.1. Grains in the ﬁlm were varied

in sizes (0.2 um or smaller). A small amount of voids was also found in the ﬁlm.

Because of the densely packed grains and the various grain sizes in the ﬁlm, the
ﬁlm was of a classic Zone T ﬁlm structure. Although the deposition temperature of 793
K (0.31 Tm of aluminum nitride) could result in a Zone 11 structure of columnar grains
[Movchan, 1969; Ohring, 1991], the discontinuous deposition process postponed the
Zone II structure from forming. Hence, the nitride ﬁlms deposited at 793 K on SiC

particles had a Zone T type in ﬁlm structure.

Aluminum Oxide Film Deposited at Room Temperature The A1203 ﬁlm
deposited on both planar substrate and the SiC particles at room temperature using an
IBAD process were opaque and black in color. The ﬁlm, shown in Figure 4.5 (a), had a
relatively smooth and dense morphology on its SiC particulate substrate. A more detailed
SEM photo with a higher magniﬁcation, shown in Figure 4.5 (b), indicates that the ﬁlm

also had globular features.

As compared with the nitride ﬁlms deposited at 593 K, the globular clusters in the
oxide ﬁlms were smaller and averaging 0.15 pm in size. The grains in the clusters were
also smaller (less than 0.1 um). Fewer voids in between the clusters were shown as
compared with the nitride ﬁlm deposited at 593 K as well. The globular feature was

more evident in the oxide ﬁlms covering areas with steps on SiC particles. The ﬁlm

47

structure of the oxide ﬁlm was classiﬁed as the Zone T structure [Movchan, 1969], with

some globular features.

Discussion on the Globular Structure in Nitride Films The globular
structure in the nitride ﬁlm deposited on SiC particles at 593 K was quite unusual. Such
a feature could be resulted from a combined inﬂuence of the deposition temperature and

the constant levitation of the particulate substrates.

At a deposition temperature of 593 K (0.24 Tm of aluminum nitride), the enhanced
atomic mobility promoted both continued nucleation of grains during deposition and
subsequent grain growth [Thomton, 1977 and 1989]. With a continuous deposition such
as in the situation of deposition on planar substrates, the resultant ﬁlms consisted of both
small and large grains, as well as of a dense structure (the Zone T Structure). 0n
particulate substrates, however, the ﬁlms were formed discontinuously because the

constant levitation of the substrates interrupted the ﬁlm growth.

It was observed in this study that, at an early stage of the ﬁlm formation on SiC
particles at 593 K, individual crystallites and clusters of crystallites were formed on the
substrate surface. Figure 4.6 (a) shows the morphology of a SiC particle after being
coated with aluminum nitride at 593 K for 2 hours. A magniﬁed SEM picture, Figure 4.2
(b) indicates that the aluminum nitride had not formed a continuous ﬁlm on the SiC
particle after the 2-hour deposition period yet. The individual crystallites and clusters of
crystallites had a size of 0.3 pm or smaller.

The development of the globular structure in the nitride ﬁlms deposited at 593 K
is speculated as a result of following steps. The deposition temperature was in favor of

large crystals developing on substrate surfaces when the ﬁlm was initiated. The

48

levitation of the substrates, however, stopped the further deposition necessary to form a
continuous ﬁlm on the substrate, leaving islands of large crystallites on the substrate

surface. These islands of crystallites increased the degree of roughness on the surface.

Because of the levitation effect, the incident angle of the depositing ﬂux changed
constantly when the particles were exposed to the deposition ﬂux. Combining the
constant changing in the incident angle and the increasing in surface roughness, the
extruding crystallite islands eventually grew into globular grain clusters on the substrate
surfaces. The globular clusters stopped to grow after being in contact with each other.
Voids were left at the boundaries in between globular clusters. Therefore, the formation
of thin ﬁlms on levitated particulate substrates at 593 K was a repetition of the formation

of individual islands of crystallites and the subsequent growth of the islands into spheres.

As compared with the ﬁlm deposition at 593 K, substrate surface provided more
suitable sites for ﬁlm nucleation at higher deposition temperatures (such as 793 K). The
higher temperature increased surface atorrric mobility, enabling the ﬁlm to horizontally
grow along the substrate surface. Hence the degree of self-roughening on the substrate
surface was less for ﬁlms deposited at a higher temperature regime. Consequently, ﬁlms

deposited at higher temperature regime were denser and less “globular” in ﬁlm structure.

Composition of Aluminum Oxide Films

XPSl was used to study the composition because of XPS’ good depth resolution
and fairly good compositional accuracy. The oxide ﬁlm was found to be predominately
aluminum oxide. Figure 4.7 (a) and (b) show an oxygen peak (530.3 eV) and an

aluminum peak (74.2 eV), both originating from aluminum oxide. No other chemical

 

1 The principles of the XPS technique are stated in Appendix I.

49

bonds were found by XPS. The composition of the ﬁlm was calculated to be Al33062.
The composition of the oxide ﬁlm was also measured using an EDX quantitative method.
Using a commercially pure A1203 rod as a standard, the result from the EDX method

yielded a composition of Al37063 for the deposited oxide ﬁlm.

The difference in the calculated ﬁlm composition and the stoichiometry (A1203)
might not be signiﬁcant after considering the errors associated with the measurement
techniques. For light elements such as oxygen, both techniques offered an uncertainty of
10% in quantitative calculation [Moulder, 1992; Russ, 1984]. Hence, a stoichiometric
A1203 sample could yield a measured composition result in a range between A146054 and

A1340“. Therefore, the oxide ﬁlm was considered to be nominal stoichiometry.

Composition and Constituents in Aluminum Nitride Films

The composition in the nitride ﬁlms was complex due to the employment of the
reactive IBAD process. XPS was used to study the atomic and molecular composition of
the nitride ﬁlms on planar substrates. Samples with i/a ratios of 0.42 and 0.60 were used

to reﬂect the deposition conditions for the particulate substrates.

The molecular concentration of the nitride ﬁlms was estimated by assuming that
all components in the ﬁlm were stoichiometric [Hentzell, 1985; Kingery, 1976]. In
Appendix II of this dissertation, detailed calculations of determining ﬁlm molecular

concentration are presented.

50

Nitride Film Deposited with an i/a Ratio of 0.42 With an i/a ratio of 0.42,
the nitride ﬁlm was deposited with an atomic arrival ratio of N-to-Al between 0.75 to

0.821. Elements of aluminum, nitrogen, and oxygen were found in the ﬁlm.

Below the ﬁrst 5 nm depth from the ﬁlm surface, the ﬁlm composition was nearly
constant at 55 at% aluminum, 40 at% nitrogen, and 5 at% oxygen. A 15 at% of oxygen
was found at the ﬁlm surface. Figure 4.8 shows an XPS concentration proﬁle in the near-
surface region of the ﬁlm. On the surface, the composition was measured to be 48 at%

aluminum, 37 at% nitrogen, and 15 at% oxygen.

The interior region of the ﬁlm consisted of aluminum nitride, aluminum nitrate,
and metallic aluminum. Table 4.2 lists the XPS results from the interior region of the

ﬁlm. The nitrogen XPS peak with a binding energy of 397.2 eV demonstrated a strong

Table 4.2. Composition of the Nitride Film Deposited with an i/a = 0.42,
at a Region 40 nm below the Film Surface.

 

 

 

 

 

Element Peak (eV) Bond Type at% of Total Content
73.8 Mixed “1 43
Al (Aluminum)
72.9 A1 11
0 (Oxygen) 532.7 Nitrate 5
N (Nitrogen) 397.2 Mixed [21 41

 

 

 

 

 

Film Molecular Concentration: 72.8% AlN + 26.1% Al + 1.1% Al(N03)3

 

 

[1]: A possible mixture of aluminum nitride and nitrate.
[2]: Aluminum nitride with possible nitrate.

 

1 As discussed previously, the percentage of N; existing in a nitrogen ion beam from a Kaufman ion
source was ranged from 78.7% to 96% [Netterﬁeld, 1988; Taylor, 1981]. The remaining species in the ion
beam were mostly N + . Hence an i/a ratio of 0.42 was equivalent to a N/Al ratio between 0.75 to 0.82.

51

aluminum nitride (AlN) characteristic. The oxygen XPS peak, as shown in Figure 4.9 (a),
indicated that the oxygen in the ﬁlm was in a form of aluminum nitrate (Al(NO3)3). The
aluminum peak indicated the presence of both a metallic aluminum (Al) component and a

possible mixture of nitride and nitrate, as shown in Figure 4.9 (b).

Compared with the interior region, the ﬁlm surface region had an additional
component of aluminum oxide. Table 4.3 lists the XPS results from the surface region of
the ﬁlm. As indicated in Figure 4.10 (a), the oxygen XPS peak from the ﬁlm surface
showed a component at 531.8 eV that originated from aluminum oxide (A120 3)
molecules. The aluminum XPS peak, as shown in Figure 4.10 (b), was a convoluted
result of characteristic peaks from components of metallic aluminum (Al) and a mixture

consisting probably aluminum nitride, nitrate, and oxide.

Table 4.3. Composition of the Nitride Film Deposited with and i/a = 0.42,
at the Surface Region.

 

 

 

 

 

 

 

 

 

 

 

Element Peak (eV) Bond Type at% of Total Content
74.1 Mixed [‘1 39
Al (Aluminum)

72.9 A1 9

N (Nitrogen) 397.4 Mixed [2] 37
532.5 Nitrate 7

0 (Oxygen)
531.8 A1203 8

Film Molecular Concentration:76.5% AlN + 15.9% Al + 1.7% A1(NO3)3 + 5.9% A1203

 

 

[I]: A possible mixture of aluminum niuide, oxide and nitrate.
[2]: Aluminum nitride with possible nitrate

52

Based on the calculations in Appendix II of this dissertation, the ﬁlm had a
molecular concentration of 72.8% AlN, 26.1% Al (metallic aluminum), and 1.1%
AI(NO3)3 in the ﬁlm interior. At the surface region, the calculated ﬁlm composition was
76.5% AlN, 15.9% Al, 1.7% Al(NO3)3, and 5.9% A1203. The results are listed in Table 4.2
and Table 4.3

Nitride Film Deposited with an i/a ratio of 0.60 With an i/a ratio of 0.60, the
nitride ﬁlm had a NzAl arrival ratio between 1.07 to 1.18 during deposition. The XPS
study found a ﬁlm composition of 44 at% aluminum, 43 at% nitrogen, and 13 at%
oxygen in the top surface layer of 10 nm thick. The oxygen concentration was reduced to

a level of 7 at% inside the ﬁlm, with 47 at% aluminum and 46 at% nitrogen in balance.

Table 4.4 lists the measured atomic composition and the calculated molecular

composition of the ﬁlm surface region. On the ﬁlm surface, the XPS study showed the

Table 4.4. Composition of the Nitride Film Deposited with an i/a = 0.60,

 

 

 

 

 

at the Surface Region.
Element Peak (eV) Bond Type at% of Total Content
531.7 A1203 6
0 (Oxygen) 533.3 A1(NO3)3 7
N (Nitrogen) 397.2 Mixed [1] 43
A1 (Aluminum) 73.9 Mixed [21 44

 

 

 

 

 

Film Molecular Concentration: 94.3% AlN + 3.4% A1203 + 2.3% Al(NO3)3

 

[1]: Aluminum nitride with possible nitrate.
[2]: A possible mixture of aluminum nitride, oxide, and nitrate.

53

 

dominant component of aluminum nitride (AlN). The oxygen XPS peak indicated the
existence of both aluminum oxide (A1203) and aluminum nitrate (A1(N03)3), as shown in
Figure 4.11. The calculated ﬁlm molecular composition, as described in Appendix II,

was 2.3% A1(N03)3, 3.4% A1203, and 94.3% AlN.

In the interior region 20 nm below the ﬁlm surface, aluminum nitride (AlN) and
aluminum nitrate (A1(N03)3) were found. The XPS peak of oxygen in this area, as shown

in Figure 4.12, indicated the existence of an aluminum nitrate (A1(N03)3) component.

To calculate the molecular composition in the ﬁlm, a small amount of metallic
aluminum (Al) was also assumed to exist in the ﬁlm, even though the XPS deconvolution
method did not separate the Al component from the dominant aluminum peak at 74.1 eV.
Calculations, detailed in Appendix H, yielded the ﬁlm’s molecular composition of 92.9%
AlN, 1.7% A1(N03)3, and 5.4% Al. In atomic composition terms, the calculation indicated
that the Al component constituted 2.5 at% in the ﬁlm. Table 4.5 lists both the measured

atomic composition and the calculated molecular composition of the ﬁlm.

Table 4.5. Composition of the Nitride Film Deposited with an i/a = 0.60,
at a Region 20 nm below the Film Surface.

 

 

 

 

Element Peak (eV) Bond Type at% of Total Content
Al (Aluminum) 74.1 Mixed [‘1 47
N (Nitrogen) 397.4 Mixed [21 46
0 (Oxygen) 532.5 Nitrate 7

 

 

 

 

 

Film Molecular Concentration: 92.9% AlN + 5.4% A1 + 1.7% A1(N03)3

 

[I]: A mixture of aluminum nitride and nitrate. Small amount of metallic aluminum may also exist.
[2]: Aluminum nitride with possible of nitrate.

54

 

Origin of Oxygen inside Nitride Films The existence of oxygen in the so—call
“nitride” ﬁlms was attributed to the incorporation of oxygen atoms in the ﬁlms by the

impingement of residual oxygen and water molecules during the deposition process.

In the deposition environment, oxygen atoms resided in two types of residual gas
molecules: 02 and H20. Both molecules produced a ﬁnite impingement rate on the
growing ﬁlm surface. As discussed in Chapter 2, the molecular concentration of an
impurity due to residual gas molecule impingement was calculated by comparing the
impinging rate of the residual gas molecules with the arrival rates of coating materials
[Chambers, 1989; Ohring, 1991]. In the deposition chamber, the impinging rate of

residual gas molecules (18) was expressed as:

J, = 2.635 ><1020 (molecules- cm2 .s" ).

a!»
"I

Here, P is the partial pressure (in Pa) for the residual gas, and T is the environment
temperature (in K). Mg is the molar weight of gas molecules, which are 32 g/mole for
oxygen and 18 g/mole for water. Hence, at an environment temperature of 300 K and a
partial pressure of 1.3x10'5 Pa, the impinging rate of residual oxygen molecules was
calculated to be 3.5x 10l3 molecules/cmzlsec. For residual H2O molecules, an impinging

rate of 4.7x 10l3 molecules/cmZ/sec was estimated with the same conditions.

The water molecular impinging rate was equivalent to the oxygen atomic
impinging rate contributed by water molecules. The oxygen atomic impinging rate from
oxygen molecules was equivalent to twice the molecular impinging rate given by the

above equation.

Assuming the oxygen atomic concentration in the ﬁlm was low and all the
impinging materials (including aluminum, nitrogen, and oxygen) were incorporated into

the growing ﬁlm, the oxygen atomic concentration (Co) in the ﬁlm was then calculated as

55

A0 A0 A0 J o

C = z = = .
0 AA,+AN+A0 AM+AN A ”ii I ”14’-
” A ” J

N N

 

 

 

Here, J0, 1,11, and JN are the arrival rate of oxygen, aluminum, and nitrogen atoms,
respectively. AM, AN, and A0 are the number of aluminum, nitrogen, and oxygen atoms in

the ﬁlm, respectively.

In the reactive IBAD used for nitride ﬁlm depositions, the arrival rate of nitrogen
atoms (JN) was proportional to the ion beam density (1,) received on the ﬁlm surface from
the assist-nitrogen ion beam. According to Netterﬁeld [1988], J N = 1.7871,. Because
J/JA; is the same as the i/a ratio, hence, the oxygen atomic concentration in the ﬁlm due

to the impingement of the water vapor molecules was rewritten as:

Puzo

P

= 3.47 x 10‘9 ”=0

1,. 1+_’_~_ ﬁ J, ”—1.— Jr
1.7871,. 1.787-t/a

The oxygen atomic concentration due to the residual oxygen molecules was then

Cg=0 = 3.47 x10”

 

 

P02

 

 

P
C32 = 5.21 x 10'9 j = 5.21x10'9
J, 1+——"— JT J, 1+—1—.— JT
1.787], 1.787 - t/a

Therefore, the total oxygen atomic concentration in the ﬁlm was then
1

1,.(1 ———1——]~/? .

+
1.787 - i/a

 

C0 = C320 + cg: = (3.4710,,20 + 5.21P02 )x10‘9

During the nitride ﬁlm deposition with an i/a ratio of 0.42, the residual partial
pressures for both oxygen (P02) and water ( P1120) were about 1.3x10'5 Pa. The ion beam
density (J,-) was about 3.4x10l4 (ions- cm’2 -s’1 ). The environment temperature (T) was
350 K. Therefore, the total oxygen concentration in the ﬁlm was estimated to be 7.6 at%.

The measured oxygen concentration was 5 at% in the interior of the ﬁlm.

56

For the ﬁlm deposited at 593 K with an i/a ratio of 0.6, the ion beam density (1,)
during the deposition was 5.2x 1014 (ions- cm‘2 -s" ), and the environment temperature (T)
was at 300 K. The residual partial pressures for both oxygen and water molecules were
1.3x 10'5 Pa. Based on the above conditions, the oxygen concentration in the ﬁlm was

estimated to be 6.0 at%. The measured value was 7 at%.

Hence, the oxygen content in the nitride ﬁlms probably originated from the
impingement of the residual oxygen and water molecules during the ﬁlm depositions.
The results reafﬁrmed the notion that even high vacuum conditions may lead to a

signiﬁcant contamination when deposition rates are low [Ohring, 1991].

Discussion of Nitride Film Deposition Process on Planar Substrates Based
on the composition study, nitride ﬁlms were actually formed in two stages of Active
Deposition and Oxidization. In the ﬁrst stage of Active Deposition, the atomic
concentrations of aluminum, nitrogen, and oxygen in the ﬁlm were determined by the
relative atomic arrival rates during the deposition process. Originating from the
impingement of residual oxygen and water molecules during deposition, oxygen atoms
were bound with nitrogen and aluminum atoms as aluminum nitrate (A1(N03)3). Under
the deposition conditions used in this study, small amount of metallic aluminum was left
in the ﬁlm after the ﬁrst deposition stage. The dominant component in the ﬁlm was

aluminum nitride (AlN).

During the second stage of Oxidization, all or part of the metallic aluminum
component was further oxidized upon exposure to air after deposition. Because of the
dense structure of the ﬁlm, the second stage of the ﬁlm formation only affected the top

surface region of the ﬁlm. The oxidized depth of the ﬁlm was about 5 to 10 nm.

57

With different i/a ratios, deposited nitride ﬁlms differed in the amount of metallic
aluminum content inside the ﬁlms. At higher i/a ratios, more aluminum atoms were
bonded with nitrogen forming aluminum nitride (AlN). For instance, in terms of
molecular concentration, the nitride ﬁlm with an i/a ratio of 0.60 consisted of 5.4%
metallic aluminum, whereas 26.1% of metallic aluminum was left in the ﬁlm deposited
with an i/a ratio of 0.42. On the surface, the ﬁlm with a lower i/a ratio formed more
aluminum oxide (A1203) component by oxidation. Some of the metallic aluminum

component on the surface might even remained.

Discussion of Composition of Nitride Films on SiC Particles Because SiC
particles were levitated during deposition, the growth of the nitride ﬁlms on SiC particles
was of an intermittent nature. The resultant ﬁlm on SiC particles experienced a multiple
alternation of the two deposition stages of Active Deposition and Oxidization. Each
second stage of oxidization was occurred in the deposition environment by the

impingement of residual oxygen and water molecules.

Each ﬁrst stage (the Active Deposition stage) of the ﬁlm deposition on SiC
particles actually resembles to the ﬁlm deposition process on planar substrates in terms of
deposition rate and conditions. The duration of each active deposition stage on SiC
particles, which was the same as the time of a particle exposing itself to the depositing
ﬂuxes each time, was observed in this study to be no more than 10 seconds. Hence, the
thickness of the ﬁlm laid down at each ﬁrst stage was no more than 2 nm (based on the

measured deposition rate on planar substrates).

After each active deposition stage, the particle was levitated away from the
incoming depositing ﬂuxes. Because of the continuation of the impingement of residual

oxygen and water molecules at the absence of the incoming aluminum ﬂux, the freshly

58

deposited ﬁlm layer underwent the second stage (the Oxidization stage) of ﬁlm formation

inside the deposition environment.

As observed in this study, the second stage of the ﬁlm formation took place at the
top surface region of 5 to 10 nm thick of the ﬁlms. Because the thickness of the freshly
deposited ﬁlm layer after each ﬁrst stage was no more than 2 nm on SiC particles, the
subsequent oxidization stage was able to take place in the entire thickness of the freshly
deposited ﬁlm layer. Under current deposition conditions, the impingement of residual
oxygen and water molecules during each oxidization stage supplied enough oxygen
atoms to oxidize the entire metallic aluminum component in the freshly deposited ﬁlm
layer. Hence, the resultant nitride ﬁlm on particles should consist of aluminum nitride,

aluminum nitrate, and aluminum oxide.

To estimate the nitride ﬁlm atomic composition and molecular constituents, it was
assumed that the N-to-Al atomic ratio in the ﬁlms was the same as the ratio of the arriving
coating ﬂuxes. The oxygen content in the ﬁlm after each ﬁrst stage of the ﬁlm formation
was presumably determined by the impingement of residual oxygen and water molecules.
After the each ﬁrst stage of the ﬁlm formation, metallic aluminum component in the ﬁlm

was presumably oxidized into aluminum oxide in the subsequent oxidation stage.

For the conditions of the 593 K deposition of nitride ﬁlm on SiC particles, an
averaged i/a ratio of 0.56 gave an N—to-Al ratio of unity [Netterﬁeld, 1988]. Using above
assumptions, the ﬁlm atomic composition was calculated to be 45.4 at% aluminum, 45.4
at% nitrogen, and 9.2 at% oxygen. In terms of molecular composition, the ﬁlm consisted
of 1.8% aluminum nitrate (A1(N03)3), 1.8% aluminum oxide (A1203) and 96.4%
aluminum nitride (AlN). (Appendix II of this paper presents detailed calculations of the

estimation of the nitride ﬁlm composition on the SiC particles.)

59

Similar estimations were also made for the nitride ﬁlms deposited on the SiC
particles at 793 K with an i/a ratio of 0.53. It was estimated that the ﬁlm had an atomic
composition of 45.7 at% aluminum, 43.3 at% nitrogen, and 11.0 at% oxygen. The

molecular composition of the ﬁlm was 94.4% AlN, 1.5% A1(N03)3, and 4.1% A1203.

The above estimations were based on a N-to-Al ratio converted from the i/a ratio
by using a factor of 1.787 [Netterﬁeld, 1988]. Taylor [1981] suggested a higher
converting factor of 1.961. Based on Taylor’s factor, a higher nitrogen atomic
concentration was estimated in the ﬁlm, which resulted in a higher concentration of AlN
in the ﬁlm with less A1(N03)3 and A1203. Using Taylor’s data, ﬁlms under our
experiment conditions yielded more than 98% of A1N with a balance of A1(N03)3 only.
However, the measured results from this study indicated that the ﬁlms consisted of oxide

content. Therefore, the estimations based on Netterﬁled’s data were more reliable.

In conclusion, nitride ﬁlms deposited on SiC particles had a dominant aluminum
nitride molecular content. The molecular concentration of aluminum nitride was

estimated between 94% to 96%, with aluminum nitrate and aluminum oxide in balance.

Summary

With IBAD techniques, averaged 0.5 pm thick aluminum oxide and nitride ﬁlms
were deposited on SiC particulate substrates. 0n SiC particles, the oxide ﬁlm had a Zone
T structure with some globular feature. The grains were 0.1 pm or smaller. The nitride

ﬁlm deposited on SiC particles at 593 K had a voided globular ﬁlm structure with

 

I The difference in conversion factor between Netterﬁeld’s and Taylor’s study is due to the difference in

their N; -to- N + ratio measurement in the ion beam from a Kaufman ion source. Netterﬁeld’s study

[1988] reported a 78.7-to-21.3 ratio, whereas Taylor [1981] reported a 96-to-4 ratio. Hence, the ratio
between nitrogen atom and ion density in the ion beam (the conversion factor) is 1.787 based on the
Netterﬁeld’s result. Based on Taylor’s measurement, the ratio is 1.96.

60

clusters of small grains. The structure for the nitride ﬁlm was a mixture of Zone I and
Zone T types. The grains in the ﬁlm were less than 0.2 pm in size, and grouped into
clusters sizing of 0.3 pm in average. The nitride ﬁlms deposited at 793 K had a dense
Zone T structure with a grain size of 0.2 pm or smaller. The composition of the oxide
ﬁlm was nominally stoichiometric. The nitride ﬁlms on SiC particles were estimated of a
N/Al ratio similar to the N/Al arrival ratio in the coating ﬂuxes, and an oxygen
concentration near 10 at%. Within the “nitride” ﬁlms, aluminum nitride constituted

about 95% of ﬁlm molecules, balanced with aluminum nitrate and aluminum oxide.

2. Thin Film Thermomechanical Behavior: Annealing Tests

It was crucial for thin diffusion barrier ﬁlms to maintain a full coverage on their
substrates during applications. An understanding of ﬁlm structural evolution helped to
study the viability and the effectiveness of the thin diffusion barrier ﬁlms. Annealing

tests were designed to reveal the thermomechanical behaviors of the ﬁlms.

Annealing at 1273 K in Air: Nitride Films on SiC Particles

SiC particles with aluminum nitride coatings were heated in air to 1273 K, and
subsequently air-cooled to room temperature. Both nitride ﬁlms (deposited at 593 K and
793 K) showed adequate adhesion to their SiC particle substrates during the test. Neither
ﬁlm showed any signiﬁcant change in morphology in the process, as compared with the
ﬁlms’ original states. Figure 4.13 shows the morphology of the ﬁlm deposited at 793 K

after the test. No cracks were found in the ﬁlms under magniﬁcations up to 7000x.

61

The nitride ﬁlms deposited on 593 K SiC particles also survived after being
quenched in liquid nitrogen (77 K) from 1273 K. No ﬁlm delamination from the SiC
particulate substrates was found. Even though some of the SiC particles themselves were
cracked during the process, the ﬁlms on these cracked SiC particles remained attaching to

their substrates, as shown in Figure 4.14.

Annealing at 1273 K in Air: Oxide Films on SiC Particles

As a comparison, the A1203 ﬁlms deposited on SiC particles were studied at the
conditions of annealing at 1273 K in air and subsequent air-cooling to room temperature.

Numerous cracks were found in the ﬁlm after the process, as shown in Figure 4.15.

At 1273 K (0.55 Tm for A1203), the annealing temperature facilitated the grain
growth in the ﬁlm. As shown in Figure 4.15, the average grain size in the oxide ﬁlm
increased from their as-deposited size (as shown in Figure 4.5) after annealing. The
grains in the oxide ﬁlm after annealing had a size comparable to the ﬁlm thickness
averaged at 0.5 pm. Consequently, a large amount of grain boundaries extended from the
ﬁlm surface to the SiC substrate. Furthermore, an agglomeration of pores at the grain

boundary region [Kingery, 1976] widened the boundaries, resulting in micro-cracks.

As a diffusion barrier to nickel aluminide matrices, A1203 ﬁlms on SiC particles
were usually subjected to a consolidation process with a temperature of 1373 K or higher.
As indicated in this experiment, the oxide ﬁlms could not retain a full coverage on SiC
substrates in such a consolidation condition. For this reason, the A1203 ﬁlm deposited on
SiC particles was not a suitable diffusion barrier material in SiC particle reinforced nickel

aluminide matrix composites.

62

Summary of Annealing Tests in Air

The annealing tests demonstrated that the oxide ﬁlm had little tolerance for a
moderate annealing temperature (1237 K). The aluminum nitride ﬁlms were capable of
retaining adhesion and a full coverage on SiC particulate substrates. The results indicated
that the oxide ﬁlm deposited on SiC particles was not a practical candidate as a diffusion
barrier in SiC particle reinforced nickel-bearing matrix. The nitride ﬁlms, on the other
hand, showed their potential as effective barriers in the composites. The nitride ﬁlms
were studied further, as presented next, with respect to ﬁlm structural stability under the

annealing conditions of composite consolidations employed in this research.

Annealing in Vacuum: Nitride Films Deposited at 793 K

Films Annealed at 1373 K for One Hour This experiment simulated the
composite consolidation conditions for SiC particle reinforced y-Ni3Al composites. In

vacuum, the coated SiC particles were annealed for one hour at 1373 K.

Figure 4.16 shows SEM images of ﬁlm morphology on an annealed particle. The
aluminum nitride ﬁlm retained adhesion to the SiC substrate. Neither cracks nor visible
voids were found on the surface of the ﬁlm under a magniﬁcation of 20,000x. There
were no obvious changes in ﬁlm morphology in the annealed ﬁlms as comparing with the

as-deposited ﬁlms as shown in Figure 4.4.

Films Annealed at I 5 73 K for Four Hours The nitride ﬁlms deposited on
SiC particles at 793 K were annealed in vacuum at 1573 K for four hours. Although

1573 K was not used to consolidate the composites in this study, the annealing condition

63

was studied because it was the highest temperature at which all the nitride ﬁlms retained
a full coverage on the SiC substrates. At temperatures above 1573 K, cracks were found

in some of the ﬁlms.

Figure 4.17 shows the morphology of the ﬁlm after the annealing. The grain
boundaries in the ﬁlm and voids in between grains became visible, as compared with the
ﬁlm morphology after being annealed at 1373 K. Besides a few large globular-structured
grain clusters on the ﬁlm surface, the grains in the ﬁlm were also different in size. Some
of the grains were about 0.2 pm in size, and were surrounded by smaller grains. Such a
mixing of different grain sizes was a classical phenomenon in grain growth process in

which the larger grains enlarged at the expense of the smaller ones [Kingery, 1976].

The appearance of visible voids on the ﬁlm surface after the annealing indicated
the progress of pore agglomeration at grain boundaries accompanied the grain growth

[Kingery, 1976].

Films Annealed at 1673 K for Four Hours This annealing condition was
used for the consolidation of B-NiAl matrix composites in this research. After annealing,
the nitride ﬁlms deposited at 793 K lost full coverage over the SiC particulate substrates.
Small patches of SiC substrate surface were observed in SEM images. Figure 4.18(a) is a
SEM image showing an annealed SiC particle with the nitride ﬁlm. Figure 4.18(b) and

(c) show more details of the morphology of the ﬁlm at higher magniﬁcations.

Most of grains in the annealed ﬁlm ranged from 0.3 to 0.6 pm in size, which was
comparable to the ﬁlm thickness of 0.5 pm in average. Smaller spherical particles of less
than 0.1 pm in size were found surrounding the bigger ones, as shown in Figure 4.18(c).

As discussed in the section of ﬁlm composition, the nitride ﬁlm consisted of about 5%

secondary phases (aluminum oxide and aluminum nitrate). The smaller spherical

particles around the bigger grains were presumably those secondary phases.

Because the grain size in the nitride ﬁlm was comparable to the ﬁlm thickness
after the annealing at 1673 K, the grain boundaries in the ﬁlm extended from the ﬁlm
surface to the SiC substrates. During grain growth process, pores in the ﬁlm traveled
along with the grain boundaries, and agglomerated there [Kingery, 1976]. The

agglomerated pores widened the grain boundaries, resulting in micro-cracks in the ﬁlm.

This experiment suggested that the aluminum nitride ﬁlms deposited on SiC
particles at 793 K would not be able to maintain an effective coverage on the SiC
substrates during a 1673 K annealing process for 4 hours. Because this annealing
condition was necessary to consolidate the B-NiAl matrix composite, therefore, the
aluminum nitride ﬁlm deposited at 793 K would not be able to protect the SiC particles

from chemical reactions with the B-NiAl matrix during the consolidation process.

Annealing in Vacuum: Nitride Films Deposited at 593 K

As compared with nitride ﬁlm deposited at 793 K, aluminum nitride ﬁlms
deposited on SiC particles at 593 K had a different ﬁlm structure after deposition. The
as-deposited ﬁlms of 593 K consisted of globular clusters of grains, as comparing with
the ﬁlms of 793 K with a smooth and dense structure. For the annealing tests in vacuum,

two temperatures (1573 K and 1673 K) were applied to the ﬁlm deposited at 593 K.

Films Annealed at I 5 73 K for Four Hours The annealed nitride ﬁlms on SiC

particles showed a globular structure of grain clusters, similar to the structure of a non-

65

annealed ﬁlm. Figure 4.19 shows an annealed SiC particle with the nitride coating and
the ﬁlm morphology on the coated particle. The nitride ﬁlm retained the coverage over

the SiC particle.

Compared with the as-deposited ﬁlm shown in Figure 4.3, boundaries between
individual grain clusters in the annealed ﬁlm were better deﬁned, and voids in between
the clusters were more visible. The agglomeration of pores at the cluster boundaries
indicated grain growth in the ﬁlms during the annealing process [Kingery, 1976].
Because there was no obvious difference between the cluster sizes before and after the
annealing, as indicated in Figure 4.3 and 4.19, the grain growth process likely occurred

within each individual grain clusters.

Films Annealed at 1673 K for Four Hours The SiC particles were entirely
covered with the nitride ﬁlm after the annealing at 1673 K for 4 hours. Figure 4.20

shows a coated SiC particle and the ﬁlm after the annealing.

The structure of the annealed ﬁlm was classiﬁed as globular, with an averaged
grain cluster size of 0.4 - 0.5 pm. Comparing images in Figure 4.19 and 4.20, the 1673 K
annealing produced a denser ﬁlm than the 1573 K annealing. The annealed ﬁlm

appeared slightly denser than the as-deposited ﬁlms (shown in Figure 4.3) as well.

Grains in individual grain clusters coalesced into larger ones as compared with the
original grain sizes. Some of the grains in the clusters were 0.3 - 0.4 um in size. Much

smaller grains (less than 0.1 um) were found surrounding the large ones.

The grain coalescence (growth) was a process in which some grains increased
their size at the expense of their neighboring grains. Because of the voided globular

structure of the ﬁlm, the grain growth process was primarily developed within each grain

66

cluster. As a result, the size of the large grains in the ﬁlm approximately equaled to the
original grain cluster size of 0.3 pm. The smaller grains of a size of 0.1 pm or less were

presumably the secondary phases (aluminum oxide and aluminum nitrate).

The denser structure in the annealed ﬁlm was attributed to a solid-state sintering
process occurring in the ﬁlm during the 1673 K annealing. Sintering process densiﬁed
the ﬁlm by eliminating the original voids in the ﬁlm [Kingery, 1976]. However, the
primary globular structure of the annealed ﬁlm indicated that the densiﬁcation via solid-

state sintering was not completed.

Therefore, the original globular structure of the nitride ﬁlm did not change after
the annealing process at 1673 K. The grain growth process was primarily limited within
the original grain clusters, resulting in a grain size comparable to the original cluster size.
Limited solid-state sintering process during the annealing helped to densify the originally
voided ﬁlm. The resultant ﬁlms after the annealing retained their coverage on the SiC

particulate substrates.

Discussion of Nitride Film Annealing in Vacuum

The vacuum annealing tests of the nitride ﬁlms strongly indicated the inﬂuence of
the original ﬁlm structure on the thermomechanical behavior of the ﬁlms. At elevated
temperatures, the extent of grain growth in a thin ﬁlm determined the survivability of the
integrity of the ﬁlm. Grain growth hinderer such as voids in the ﬁlm played a profound

role in ﬁlm’s thermomechanical behavior.

As an inevitable high—temperature process, grain growth increased average grain
size in materials by combining several smaller grains into a larger one [Kingery, 1976].

Without a grain growth hinderer such as voids, ﬁlms with small grains experienced an

67

uninhibited, homogeneous grain growth in the ﬁlm at elevated temperatures. Such a
grain growth in the ﬁlm eventually led to a grain size comparable to the ﬁlm thickness,
extending the grain boundaries in the ﬁlm from the ﬁlm surface to the substrate.
Agglomerating of pores at the grain boundaries during grain growth further compromised
the integrity of the annealed ﬁlm. The nitride ﬁlm deposited on SiC particles at 793 K

was such an example.

By being a voided structure, ﬁlms with an originally globular structure not only
provided the grain growth hinderer such as voids, but also set a stage for solid-state
sintering taking place at high temperatures [Kingery, 1976]. The original voided globular
structure conﬁned the grain growth process primarily within individual grain clusters.
The sintering process densiﬁed the ﬁlm to a certain extent. Because voids in ﬁlm stalled
the grain growth process [Kingery, 1976], the progress of grain coalescence in the
originally voided ﬁlm was less advanced than that in a ﬁlm with an original dense
structure. The grain size in the ﬁlm deposited at 593 K was found to be less than that in

the ﬁlm deposited 793 K after the same annealing.

Therefore, the original voided globular structure was actually preferred in this
study for thin ﬁlms under high-temperature (>0.5Tm of the ﬁlm material) applications, as
comparing to the originally dense ﬁlm structure with small grains. In this study, the
nitride ﬁlm deposited on SiC particulate substrate at 593 K consisted of a globular ﬁlm
structure with clusters of ﬁne grains. This nitride ﬁlm was preferred as diffusion barrier

at high temperatures to nitride ﬁlms deposited at 793 K with a dense structure.

68

3. Ni Diffusion in Aluminum Nitride, a Cursory Study

There was limited information available regarding nickel diffusion in aluminum
nitride. A cursory study was conducted to determine if Ni has low lattice diffusivity in
the nitride at temperatures applied in this research. The extent of Ni diffusion in both
monolithic aluminum nitride plate and deposited aluminum nitride ﬁlm were studied. In

this investigation, only the extent of Ni diffusion in the nitride was studied.

Ni Diffusion at the Interface of Aluminum Nitride and Nickel Aluminide Plates

An aluminum nitride plate was diffusion bonded in vacuum with a y-N 13A] plate
at 1473 K for 32 hours. After revealing the interface of the diffusion couple, an EDX Ni
line-proﬁle showed that there was no signiﬁcant lattice diffusion of Ni across the

interface. The interfacial bond strength between these two plates was very weak.

Assuming a constant diffusant source situation‘, the extent of Ni diffusion in
aluminum nitride was estimated by measuring the width of Ni concentrational transition
at the interface on SEM/EDX photos. Measured from an EDX Ni line-proﬁle shown in
Figure 4.2l(a), the Ni diffusion extent was less than 1 pm (which is the spatial resolution
for the EDX technique [Russ, 1984]). An upper-bound lattice diffusivity for Ni diffusing
in the alunrinum nitride lattice was therefore estimated to be less than 2.2x 10"M cmZ/sec

under the conditions of the diffusion bonding. Accordingly, for an one-hour diffusion at

 

l A constant diffusant source situation represents an ideal diffusion model where the diffusant
concentration remains constant at the interface of a diffusion couple, and the same as in the source
material. This model is best ﬁt in a diffusion couple situation that diffusant has a high diffusivity in the
diffusant source comparing with the diffusivity in the diffusion matrix material.

69

1473 K, an upper-bound distance of Ni diffusing in the nitride via lattice diffusion was

estimated to be 0.18 uml.

A diffusion couple of planar aluminum nitride and B-NiAl was subjected to a
bonding treatment at 1673 K for 12 hours. As shown in Figure 4.21(b), the transition
distance of Ni concentration at the interface was also less than 1 pm. Hence, an upper-
bound diffusivity of Ni diffusing in the B-NiAl lattice at this temperature was estimated
to be less than 5.8x10‘l4 cmz/sec. An upper-bound lattice diffusion distance was

estimated to be 0.58 pm under a hot-pressing condition of 1673 K for 4 hours.

Ni DiffusionalPenetration in a Deposited Aluminum Nitride Film

This experiment was designed to measure the extent of Ni diffusion in a deposited
aluminum nitride ﬁlm. A y-Ni3Al plate was chosen as the source for Ni. After the nitride
ﬁlm coated plate was annealed in vacuum at 1373 K (the consolidation temperature for y-
Ni3Al matrix composites) for 30 rrrinutes, the extent of Ni diffusion in the nitride ﬁlm

from the y-Ni3Al substrate was studied by XPS.

Figure 4.22 shows a Ni concentration proﬁle in the aluminum nitride ﬁlm. In the
proﬁle, a concentration ratio of Ni-to-Al was used for the y-axis. A 6/5 power of the
distance from the ﬁlm surface was used for the x-axis, for which was proportional to the

diffusant concentration due to grain-boundary diffusion [LeClaire, 1961].

 

1 The maximum distance (x)of a diffusant can be expressed as:
x = 2 D,t
Here, the D, is the lattice diffusivity, and t is the diffusion duration. For diffusion at 1473 K between
aluminum nitride and y-Ni3Al, x = 1 um, I = 32 hours. Hence, D, was calculated to be 2.2x10'" cmzlsec.

For an one-hour diffusion duration, the upper-bound distance of Ni lattice diffusion in the nitride is then
0.18 pm.

70

As indicated in the proﬁle, the concentration of Ni content reduced sharply within
a region about 10 nm wide in the nitride ﬁlm at the ﬁlm-substrate interface. Following
that region, the concentration decreased gradually and penetrated into the ﬁlm further.
No Ni content was found on the ﬁlm surface. The total distance of Ni diffusing in the
nitride ﬁlm after the 0.5 hour annealing duration was extrapolated to be about 200 nm

(0.2 um) from the interface.

The transition region with high Ni concentrations at the ﬁlm-substrate interface
was interpreted as a result of the lattice diffusion of Ni in the ﬁlm. The region with a
gradually changed concentration proﬁle was most likely caused by grain-boundary
diffusion [Gilmer, 1976a]. The small grain size of the ﬁlm facilitated the extent of the Ni

diffusion in the ﬁlm.

For a condition of an one-hour annealing in vacuum at 1373 K, the extent of Ni
diffusion in the nitride ﬁlm was estimated to be roughly 0.3 pm from a y-Ni3Al source.
This estimation was based on the relationship of the diffusion distance being proportional

to the square root of diffusion duration.

Summary and Discussion

This cursory study showed, with both B-NiAl and y-Ni3Al as nickel sources,
nickel has very low lattice diffusivities in aluminum nitride. From a y-Ni3Al source, the
upper-bound lattice diffusivity of Ni in the nitride was 2.2x10'l4 cmz/sec at 1473 K. At

1673 K, the upper-bound diffusivity was 5.8x10'14 cmZ/sec from a B-NiAl source.

The study of Ni diffusion in a deposited aluminum nitride ﬁlm showed a greater

diffusional distance than the estimated distance based on the lattice diffusion. Because of

71

the grain-boundary diffusion effect, the diffusional penetration was about one order of

magnitude greater than the prediction from the lattice diffusion data.

From a y-Ni3Al source, the extent of Ni diffusing in the nitride ﬁlm at the
consolidation condition of y-Ni3Al composites (1373 K/l hour) was estimated to be about
0.3 pm. This distance was less than, but comparable to the thickness of the ﬁlms
deposited on SiC particulate substrates used in this study (As discussed in Section 1 of

this chapter, the ﬁlm thickness in this study was about 05:01 pm.)

As to the Ni diffusion in the nitride ﬁlm from a B-NiAl source, the upper-bound
lattice diffusion distance at the consolidation condition of B-NiAl composites (1673 K/4

hours) was estimated to be less than 0.58 pm.

Hence, with a nominal thickness of 0.5 pm, the deposited aluminum nitride ﬁlm
on SiC substrates was a good candidate to evaluate the chemical potential effect on the

effectiveness of diffusion barriers in different composite systems.

4. Effectiveness of Diffusion Barriers in Two Composites: SiC

Particle Reinforced B-NiAl and y-Ni3Al Composites

The chemical potentials were different in composite systems of SiC/B-NiAl and
SiC/y-Ni3Al. The interaction between SiC and B-NiAl was moderate as compared with
the interactions in a SiC/y-Ni3Al composite [Chou, 1991a&b]. The reaction zone
between SiC and y-Ni3Al was about an order of magnitude wider from a 1273 K
annealing than the reaction zone in a SiC/B-NiAl diffusion couple annealed at 1573 K.
With a higher Ni content in the matrix, the SiC/y-Ni3Al composite system offered a
higher chemical potential, and a much more challenging task for the diffusion barriers

than the B-NiAl matrix composite did.

72

In Section 2 of this chapter, three barrier ﬁlms (aluminum oxide deposited at
room temperature, aluminum nitride ﬁlm deposited at 593 K, and aluminum nitride ﬁlm
deposited at 793 K) demonstrated different thermomechanical behaviors during annealing
at temperatures for composite consolidation. This section addresses the effectiveness of
the ﬁlms as diffusion barriers in both SiC particle reinforced B-NiAl and y-Ni3A1

composites under the composites’ respective “hot-press” consolidation conditions.

The Control Composite Samples

The control composite samples, which used non-coated SiC particles, were
consolidated with both B-NiAl and y-Ni3Al matrices under the composites’ respective
consolidation conditions. The B-NiAl matrix composites were hot-pressed for 4 hours at

1673 K. The y-Ni3Al matrix composites were consolidated at 1373 K for one hour.

As shown in Figure 4.23(a), the non—coated SiC particles reacted with the B-NiAl
matrix. A 20 pm thick reaction zone was formed at the interface region between the SiC
particles and the B-NiAl matrix. As shown in Figure 4.23(b) of an EDX spectrum from
the reaction zone, the reaction product was a Ni-Al-Si ternary compound with a low Si

content that was qualitatively in agreement with the previous report by Chou [1991a].

Figure 4.24 shows the reaction zone in a y-Ni3Al matrix composite sample where
the non-coated SiC particles were used. The y-Ni3Al reacted severely with the SiC
reinforcements at the consolidation conditions. Layers of reaction products were formed
in the area previously occupied by the SiC particles. No remnant of the SiC particle was

found after the consolidation.

73

Aluminum Nitride Films Deposited at 593 K as Diffusion Barrier

In ,B-NiAl Matrix Composite After consolidated with the B-NiAl matrix at
1673 K for 4 hours, the coated SiC particles were protected from the chemical reactions

with the matrix by an average 0.5 um thick nitride ﬁlm deposited at 593 K.

More than 90% of the coated SiC particles were preserved with the nitride ﬁlm
intact. Figure 4.25(a) and (b) are SEM images of a coated SiC particle in the composite
after the consolidation. On this particle, part of the nitride ﬁlm was damaged, showing a
relatively smooth area. Only one strong Si signature was shown in an EDX spectrum
collected from the smooth area, as shown in Figure 4.25(c), which indicated that the ﬁlm
was damaged during the SEM sample preparation. The revealed SiC substrate in this

area showed no sign of the reaction products.

On fewer than 10% of the coated particles in the composite, the nitride ﬁlms were
broken during the hot-pressing process. In these regions with ﬁlm damages, the SiC
substrate was exposed to and chemically reacted with the NiAl matrix during the
consolidation process. Figure 4.26(a) shows one of the few such particles. The particle
was released by electropolishing to remove the matrix material. In Figure 4.26(b), a
magniﬁed picture of the ﬁlm damaged region shows reaction products similar to those in
the control samples in Figure 4.23. Figure 4.26(c) is an EDX spectrum from the substrate

at the ﬁlm—damaged region, indicating the presence of a ternary Ni-Al-Si compound.

In y-Ni 3A1 Matrix Composite The aluminum nitride ﬁlm deposited on SiC
particles at 593 K did not prevented the reactions in the y-Ni3Al matrix composite system.

A reaction zone was formed in the composites after the consolidation.

74

Figure 4.27 presents an optical micrograph of the reaction zone in the composite
containing the coated SiC particles. The structure of the reaction zone in the composite
was almost identical to the zone in the control samples shown in Figure 4.24. A broken
line at the outer perimeter of the reaction zones showed some evidence of the location, or

might even the remnants of the nitride ﬁlm.

Discussions The results of this experiment indicated that an average 0.5 pm
thick nitride ﬁlm deposited at 593 K was an effective diffusion barrier for SiC particles in
a B-NiAl matrix composite during the composite consolidation process at 1673 K.
However, the same nitride ﬁlm was not adequate to protect the SiC particles in a y-Ni3A1

matrix composite at 1373 K.

This experiment demonstrated the importance of the chemical potential in the
composite system in determining the effectiveness of the diffusion barrier. Although the
nitride ﬁlm was an effective diffusion barrier in the B-NiAl matrix composite with a
modest chemical potential, the ﬁlm did not offer any sign of stalling the reactions in the

y-Ni3Al matrix composite with a higher chemical potential.

The consolidation process for B-NiAl at 1673 K probably had an additional
advantage for the nitride ﬁlm deposited on SiC particles at 593 K. As discussed in the
section of Thermomechanical Behavior, the ﬁlm underwent a limited solid-state sintering
process on the SiC particles at 1673 K. The ﬁlm structure after the sintering process was
denser than the non-annealed state. As for the ﬁlm after a 1373 K consolidation process
for y—Ni3Al, the annealed ﬁlm was as voided as the ﬁlm in the original state. A voided

ﬁlm offered more fast-diffusion baths than a densely packed thin ﬁlm.

75

Aluminum Nitride Films Deposited at 793 K as Diffusion Barrier

In ,B-NiAl Matrix Composite SiC particles used in the composite were coated
with nitride ﬁlm at 793 K. The average ﬁlm thickness was 0.5 pm. After consolidation
with the B-NiAl matrix at 1673 K for 4 hours, the coated SiC particles were found to
react with the matrix. Figure 4.28 shows a reacted SiC particle in the B-NiAl matrix
composite. The reaction zone was similar to that in the control sample. There was no

evidence that the nitride ﬁlm had any effect in stalling the matrix-reinforcement reaction.

In 7-Ni 3A1 Matrix Composite The aluminum nitride ﬁlm deposited at 793 K
did not prevented the reactions in the SiC/y-Ni3A1 matrix composite system. Reaction
zones were formed in the composites after the consolidation. Figure 4.29 shows the
reaction zone at a location formerly occupied by a coated SiC particle. The composite
was fractured in tension at room temperature after the consolidation. The reactions
between the y-Ni3Al matrix and SiC particle consumed the entire SiC particle. No

remnant of the SiC particle was found after the consolidation.

Discussions The results of this test indicated that an average 0.5 urn thick
nitride ﬁlm deposited at 793 K was not adequate to protect the SiC particles in both [3-
NiAl matrix and y-Ni3Al matrix composites during the composites’ respective
consolidation processes. In the y-Ni3Al matrix composite, the high chemical potential in

the system was blamed as the cause of the ineffectiveness of the diffusion barrier.

In the B-NiAl composite, the thermomechanical behavior of the nitride ﬁlm
probably cost the ﬁlm’s effectiveness as the diffusion barrier. As discussed in the section

of Thermomechanical Behavior, the nitride ﬁlm deposited at 793 K lost integrity over the

76

SiC particulate substrates under the conditions for B-NiAl consolidation at 1673 K.
Uninhibited grain growth resulted in grain boundaries extending from the ﬁlm surface to
the substrate. The breakage of the ﬁlm’s continuity presumably led to the ﬁlm’s failure

to protect the substrates.

Aluminum Oxide Films as Diffusion Barrier

An average 0.5 pm thick oxide ﬁlm on SiC particles was not sufﬁcient to protect
the SiC particles from chemical reactions in both B-NiAl matrix and y-Ni3Al matrix
composites during the composites’ respective consolidation processes. As shown in
Figure 4.30, there was no evidence of stalling the reactions in both matrices as compared
with the control composite samples. Figure 4.30(a) shows the reaction products around
the remnant of a SiC particle in a B-NiAl matrix composite. Figure 4.30(b) presents the

reaction zones in a y-Ni3Al matrix composite.

As discussed in the section of Thermomechanical Behavior, the oxide ﬁlm
exhibited a higher tendency of grain growth at lower temperatures than the nitride ﬁlms
did. Because of a ﬁlm structure of densely packed small grains in the as-deposited oxide
ﬁlm on SiC particles, uninhibited grain growth in the ﬁlm during annealing eventually
cost the ﬁlm’s continuity. Numerous cracks were observed in the oxide ﬁlm after the
coated SiC particles had been annealed at 1273 K (Figure 4.15). It was not a surprise that
the oxide ﬁlm failed to protect the SiC substrates during the composite consolidation

processes at temperature of 1373 K or above.

77

Summary

Although SiC was chemically ready to react with B-NiAl at temperatures above
1573 K [Chou, 1991a], the imposition of an aluminum nitride diffusion barrier could stall
the reaction long enough for the success of consolidating the composite at 1673 K. With
an original globular ﬁlm structure and an average thickness of 0.5 pm, the nitride ﬁlm
deposited on SiC particles at 593 K was an effective diffusion barrier in SiC particle

reinforced B-NiAl matrix composites processed at a temperature of 1673 K.

This study indicated that the chemical potential in the composite system played an
important role in determining the effectiveness of the diffusion barrier. Although the
nitride ﬁlm deposited at 593 K functioned effectively as a diffusion barrier in the SiC/B-
NiAl composite system, the same ﬁlm failed to protect the SiC particles in a higher
chemical potential composite system (the SiC/y-Ni3Al composite system) during the

composite’s consolidation process at 1373 K.

The thermomechanical behavior of the barrier ﬁlms was also important. Both
densely packed ﬁne- grain ﬁlms, the oxide ﬁlm deposited at room temperature and the
nitride ﬁlm formed at 793 K, lost the ﬁlm integrity on the SiC particulate substrates
because of the uninhibited grain growth in the ﬁlms at elevated temperatures. None of
the densely packed ﬁne-grain ﬁlms functioned as the effective diffusion barrier in SiC

reinforced aluminide matrix composites in this study .

78

CHAPTER 5. CONCLUSIONS

In this study, aluminum nitride and aluminum oxide ﬁlms were used as diffusion
barriers in SiC particle reinforced nickel alurrrinide matrix composites (both SiC/B-NiAl
and SiC/y-Ni3Al systems). Several factors were found to affect the effectiveness of the
diffusion barrier ﬁlms. These factors included barrier ﬁlm materials, barrier ﬁlm
deposition conditions and resultant ﬁlm structures, barrier ﬁlm thermomechanical
behavior during compositing process, and chemical potential in the composite system. In
this study, a nitride ﬁlm with an average thickness of 0.5 urn demonstrated its capability
as an effective barrier to the catastrophic reinforcement-matrix reactions in SiC particle

reinforced B-NiAl matrix composites at a temperature of 1673 K.

A cursory investigation done in this study indicated that Ni has low lattice
diffusivity in aluminum nitride (AlN) at high temperatures. Only the upper-bound value
of the diffusivity of Ni in AW lattice was estimated in this study. The estimated upper-
bound diffusivity of Ni from a y-Ni3Al source was 2.2x10'l4 cm2/sec at 1473 K, and
5.8x10'l4 cmZ/sec at 1673 K from a B-NiAl source. In an aluminum nitride ﬁlm
deposited on a y—Ni3Al plate, Ni penetration was measured to be about 0.2 pm deep into
the nitride ﬁlm after annealing at 1373 K for 0.5 hour. The diffusion depth was

contributed predominantly by the grain-boundary diffusion in the ﬁne-grained ﬁlm.

In this study, the nitride and oxide ﬁlms were fabricated by ion-beam assisted
deposition (IBAD) processes. Both ﬁlms were deposited on SiC particulate substrates in
a specially designed particle levitation system. The oxide ﬁlm was formed at room
temperature with an assisted argon ion beam. The nitride ﬁlms were formed at elevated

temperatures reactively using a nitrogen ion beam.

79

By using the reactive IBAD technique, the nitride ﬁlm was formed in a two-stage
process of Active Deposition and subsequent Oxidization. After the active deposition
stage, the constituents in the nitride ﬁlm were found to be aluminum nitride (AlN),
aluminum nitrate (A1(N03)3), and metallic aluminum (Al). Because of the low deposition
rate on the particulate substrates, the entire Al component in the ﬁlm was oxidized in situ
by the impingement of residual oxygen and water molecules the second stage of ﬁlm

formation, the Oxidization stage.

While the oxide ﬁlm had a nominally stoichiometric composition and a single
A1203 component, the deposited nitride ﬁlm on SiC particles consisted of about 10 at%
oxygen content. The N-to-Al ratio in the nitride ﬁlm was estimated to be between 0.95
and 1.0. An estimated 95% of molecules in the nitride ﬁlms on SiC particles were AlN,

with A1(NO3 )3 and A1203 molecules in balance.

On SiC particulate substrates, the average thickness for both oxide and nitride
ﬁlms was about 0.5 um. Grains in both ﬁlms were smaller than 0.2 pm. However, ﬁlms

deposited at different temperatures were different in structures.

The nitride ﬁlm deposited at 593 K had a globular structure with clusters of
grains. Large amount of voids was visible in the ﬁlm. The average size of the clusters in
the ﬁlm was 0.3 pm. The nitride ﬁlm deposited at 793 K had a densely packed structure.
The structure of the oxide ﬁlm deposited at room temperature consisted of densely
packed grains, with some globular features. The difference in ﬁlm structure was
attributed to the combined effect of the levitation of the particulate substrates during

deposition and the difference in deposition temperatures.

At high temperatures, grain growth was observed in both oxide and nitride ﬁlms.
The ﬁlm material affected the rate of the grain growth. The oxide ﬁlm had a higher rate

of grain growth than the nitride ﬁlm did at elevated temperatures. At 1273 K, the grains

80

in the oxide ﬁlm on SiC particles coalesced to an average size that was comparable to the
ﬁlm thickness. The grain boundaries in the oxide ﬁlm extended from the ﬁlm surface to
the ﬁlm-substrate interface. Agglomeration of pores at the grain boundaries in the oxide
ﬁlm widened the boundaries into micro-cracks. At higher temperature such as 1673 K,
the nitride ﬁlm with densely packed small grains (deposited on SiC particles at 793 K)
experienced the similar process in which grain growth led towards larger grains and

micro-cracks in the annealed ﬁlm.

This study demonstrated the effect of ﬁlm structure on ﬁlm’s thermomechanical
behavior at high temperatures. The structure of the nitride ﬁlm deposited at 593 K was
globular with grain clusters. Because of large amount of voids in between the clusters,
grain growth in the ﬁlm was restricted within each cluster. At high temperatures, solid-
state sintering also densiﬁed the ﬁlm. As a result, the nitride ﬁlm not only retained the
coverage over the SiC particulate substrates after the annealing at 1673 K, but also

evolved into a ﬁlm with less voids.

This study showed the importance of the chemical potential in composite system
in determining the effectiveness of a diffusion barrier ﬁlm. In a composite system with a
high chemical potential (such as in the SiC/y-Ni3Al system), a diffusion barrier ﬁlm was
less effective than the same ﬁlm was in a composite system with a modest potential (such
as in a SiC/B-NiAl system). With an average thickness of 0.5 pm, the nitride ﬁlm
deposited on SiC particles at 593 K was an effective barrier for Ni diffusing in the SiC/[3-
NiAl composite during the composite consolidation process at 1673 K. However, the
same nitride ﬁlm did not prevent reactions between the SiC reinforcement and the y-

Ni3A1 matrix during the SiC/y-Ni3Al composite consolidation process at 1373 K.

In conclusion, this study indicated that several key factors other than diffusivity

were critical for the success of a diffusion barrier in a composite system. The chemical

81

potential in the composite system was the most crucial factor in determining the
effectiveness of the diffusion barrier. By affecting ﬁlm’s thermomechanical behavior at
elevated temperatures, the initial ﬁlm structure signiﬁcantly inﬂuenced the effectiveness

of the diffusion barriers in hi gh-temperature composites.

Though model composite systems were used in this study, it was clearly shown
that, with a proper ﬁlm structure, using aluminum nitride ﬁlms on the SiC particles was a

possible way to fabricate SiC particle reinforced B-NiAl matrix composite.

82

APPENDICES

83

APPENDIX I.

PRINCIPLES OF THE XPS TECHNIQUE

The technique of X-ray photoelectron spectroscopy (XPS, sometimes also called
ESCA) was applied in this work for thin ﬁlm composition study. In XPS, a sample is
illuminated with X-rays that excite photoelectrons from the surface of the sample
[Moulder, 1992; Walls, 1989]. The photoelectrons have binding energies that are
characteristic of individual atoms to which the electrons were bound. Because of the
shallow escape depth of the photoelectrons, XPS is a surface analysis tool with a depth
resolution of 0.2 nm. With a complementary ion-gun, XPS can be used for composition

proﬁling. The depth resolution for the proﬁling is poor compared with the static method.

The photoelectron energy is dependent on the precise chemical conﬁguration of
the atoms. For an element with several chemical bonds in a material, the element’s XPS
peak is a convolution of all the photoelectron energies from all the chemical bonds.
These bonds and their relative amount on a sample surface can be analyzed from the XPS
spectrum by deconvoluting the XPS peak. Hence, XPS is a tool to obtain chemical
information from a material's surface. Table A lists some of the binding energies for

three elements: Al, N, and O in some of their compounds [Moulder, 1992, Taylor, 1981].

For a sample that is homogeneous in the analysis volume, the molar concentration

of any constituent in the sample, C x, can be written as

_ Ix/Sx

x Eli/Si

84

where I x is the measured peak intensity for element x and S is the relative atomic

sensitivity factor for that peak. The values of S are based on empirical data, and are

matrix-dependent.

The XPS technique is a semi-quantitative one. It has a reasonable quantitative

accuracy (better than :I:10% with good calibration standards) among the surface analysis

techniques. The accuracy is inﬂuenced by the homogeneity of the sample surface, the

elements involved, and the degree of peak interference.

 

 

 

 

 

 

 

Table A. Standard XPS Binding Energy for Al, N, and O [Moulder, 1992;
Taylor, 1981].
Material Binding Energy (eV)
Al N 0
A1 72.8 - -
AlN 73.9 - 74.4 397.3 -
A1203 73.7 - 74.7 - 530.0 - 531.8
Nitrate Not Available 407.0 - 408.1 532.5 - 533.6

 

 

 

 

85

 

APPENDIX II.

CALCULATION OF NITRIDE FILM MOLECULAR CONCENTRATION FROM
XPS RESULTS

In this study, XPS results indicated that the as-deposited nitride ﬁlms consisted of
up to four components: aluminum nitride, aluminum nitrate, aluminum oxide, and
metallic aluminum. Because of the limitation of XPS method, elemental information in
some of the components could not be separated from each other. For instance, the
aluminum characteristic peaks from nitride and oxide (located at 74.15 i 0.25 eV and
74.2 i0.5 eV, respectively) could not be separated by a deconvolution method.
Additional assumptions were needed to calculate the molecular concentration of a sample
using XPS measurement. In following estimations, all components in the ﬁlm were

assumed at their stoichiometric state.

Case 1. Nitride Film with an i/a Ratio of 0.42 on Planar Substrate

Interior Region

At the interior region of a nitride ﬁlm deposited with an i/a ratio of 0.42, the XPS
results showed three molecular components: aluminum nitride (AlN), aluminum nitrate
(A1(N03)3), and metallic aluminum (Al). The atomic composition of the ﬁlm was found

of 54 at% aluminum, 41 at% nitrogen, and 5 at% oxygen.

86

For every 100 atoms at the interior region of the ﬁlm, let u be the number of
aluminum atoms, v for nitrogen, and w for oxygen, respectively. Thus, u+v+w = 100.
The 100 atoms formed a molecules of AW and b molecules of A1(N03)3, leaving c atoms

(molecules) of Al. Hence, the ﬁlm components were: aAlN + bAl(NO3)3 + CA].

The amount of individual atoms in the ﬁlm could be expressed as:

Aluminum Atom: a + b + c = u
Nitrogen Atom: a + 3b = v
Oxygen Atom: 9b = w
Therefore,

' _ _ w

a _ v /3
< b = W9
_ _ 2w
kc — u v + /9

 

Because u = 54, v = 41, and w = 5, the number of molecules formed from the 100
atoms at the interior region of the ﬁlm was therefore 39.33 for AlN (a), 0.56 for A1(N03)3
(b), and 14.11 for Al (c). For a total of 54.0 molecules (=a+b+c), the molecular
composition in the ﬁlm was 72.8% AlN, 1.1% A1(N03)3, and 26.1% Al.

The above estimation yielded a concentration of 14.1 at% metallic aluminum
component at the ﬁlm interior region. The measured concentration of metallic aluminum

was 11 at% by the XPS deconvolution method.

Surface Region

At the surface region of the nitride ﬁlm, four components were found: aluminum

nitride (AlN), aluminum nitrate (A1(N03)3), aluminum oxide (A1203), and metallic

87

aluminum (Al). The atomic composition was measured of 48 at% aluminum, 37 at%

nitrogen, and 15 at% oxygen.

For every 100 atoms at the surface region, let u be the number of the aluminum
atoms, v for nitrogen, and w for oxygen, respectively. Thus, u+v+w = 100. Furthermore,
let w, and W2 be the number of oxygen atoms bound in A1203 and A1(N03)3, respectively.
Hence, w, + M = w. The 100 atoms formed a molecules of AlN, b molecules of
A1(N03)3, and c molecules of A1203, leaving d atoms (molecules) of Al. Hence, the ﬁlm

components were: aAlN + bAl(N03)3 + c A1203 + dAl.

The amount of individual atoms in the ﬁlm could be expressed as:

Aluminum Atom: a + b + 2c + d = u
Nitrogen Atom: a + 3b = v

3c = wI
Oxygen Atom: 9b = w2

wl+w2=w

la=v—W%
.-w/
kd=u—v+2W%—ZW%

For every 100 atoms at the surface region of the ﬁlm, u = 48, v = 37, and w = 15.

Therefore,

 

Furthermore, it was measured by XPS method that w, = 8 and W2 = 7. The number of
molecules formed from the 100 atoms at the surface region of the ﬁlm was therefore
34.67 for AlN (a), 0.78 for A1(N03)3 (b), 2.67 for A1203 (c), and 7.22 for Al (d). With a
total of 45.34 molecules (=a+b+c+d), the above estimation yielded a molecular

composition of 76.5% AlN, 1.7% A1(N03)3, 5.9% A1203, and 15.9% Al.

88

The above estimation yielded a concentration of 7.2 at% metallic aluminum
component at the ﬁlm surface region. According to XPS deconvolution result, the atomic

concentration of metallic aluminum was 9 at%.

Case 2. Nitride Film with an i/a Ratio of 0.60 on Planar Substrate
Interior Region

In the interior region, XPS results showed that the ﬁlm consisted of components
of aluminum nitride (AlN) and aluminum nitrate (A1(N03)3). The atomic composition

was 47 at% aluminum, 46 at% nitrogen, and 7 at% oxygen.

Although XPS study did not show any metallic aluminum in the ﬁlm, the
following estimation was assumed that a small amount of metallic aluminum (Al) existed
in the ﬁlm. For every 100 atoms at the surface region, let u be the number of the
aluminum atoms, v for nitrogen, and w for oxygen, respectively. Thus, u+v+w = 100.
The 100 atoms formed a molecules of AW and b molecules of A1(N03)3, leaving c atoms

(molecules) of Al. Hence, the ﬁlm components were: aAlN + bAl(N03)3 + CA].

The amount of individual atoms in the ﬁlm could be expressed as:

Aluminum Atom: a + b + c = u
Nitrogen Atom: a + 3b = v
Oxygen Atom: 9b = w
Therefore,

ra=v—%
<b=%

=_2w
Lcuv+ 9

 

89

Because u = 47, v = 46, and w = 7, the number of molecules formed from these
100 atoms at the interior region of the ﬁlm was therefore 43.67 for AlN (a), 0.78 for
A1(N03)3 (b), 2.56 for Al (c). With a total of 47.01 molecules (=a+b+c), the above
estimation yielded a molecular composition of 92.9% AlN, 1.7% A1(N03)3, 5.4% Al.

The above estimation suggested that a 2.56 at% of metallic aluminum component
could probably exist in the ﬁlm. The justiﬁcation of this assumption was based on the
existence of aluminum oxide on the ﬁlm surface. The oxide was most likely formed by
the oxidation of the metallic aluminum component at the surface. For a XPS peak from a
component with such a low concentration, the neighboring dominant aluminum peaks

from nitride/nitrate components would easily overshadow it.

Surface Region

At the surface region, the ﬁlm consisted of three components: aluminum nitride
(AlN), aluminum nitrate (A1(N03)3), and aluminum oxide (A1203). The atomic

composition was 44 at% aluminum, 43 at% nitrogen, and 13 at% oxygen.

For every 100 atoms at the surface region, let u be the number of the aluminum
atoms, v for nitrogen, and w for oxygen, respectively. Thus, u+v+w = 100. The 100
atoms formed a molecules of AlN, b molecules of A1(N03)3, and c molecules of A1203.

Hence, the ﬁlm components were: aAlN + bAl(N03)3 + cA1203.

The amount of individual atoms in the ﬁlm could be expressed as:

Aluminum Atom: a + b + 2c = u
Nitrogen Atom: a + 3b = v
Oxygen Atom: 9b + 3c = w

90

Therefore,
=/-//
f: 3%‘3%+W12

For every 100 atoms at the surface region of the ﬁlm, u = 44, v = 43, and w = 13.

A

 

The number of molecules formed from the 100 atoms in the surface region of the ﬁlm
was therefore 40.13 for AlN (a), 0.96 for A1(N03)3 (b), and 1.46 for A1203 (c). With a
total of 42.55 molecules (=a+b+c), the above estimation yielded a molecular composition

of 94.3% AlN, 2.3% A1(N03)3, and 3.4% A1203.

The above estimation yielded an atomic concentration of oxygen in nitrate of 8.6
at%, and of 4.4 at% bonded in oxide. According to the XPS, the atomic concentration of

oxygen in nitrate was 7 at%, and 6 at% in oxide.

Also, because the aluminum oxide component was most likely formed by the
oxidation of the metallic aluminum component in the ﬁlm, the amount of the oxide
component at the surface region required a molecular concentration of 6.6% metallic
aluminum before the oxidation occurred. It was in a good agreement with the estimated

result of 5.4% for the interior region.

Case 3. Nitride Film Deposited on SiC Particles

General Assumptions

For nitride ﬁlms deposited on SiC particles, the ﬁlms were presumably formed in
two stages. The ﬁrst stage was the Active Deposition where most of the ﬁlm components

were deposited from the coating ﬂuxes. In the second stage, other components were

91

formed by oxidization. These two stages of deposition took place in turn during the ﬁlm

depositions because of the levitation of the SiC particles.

At the ﬁrst stage where active depositions took place, all arrival aluminum and
nitrogen atoms were incorporated into the ﬁlm on the particle surface. Such an
assumption gave the ﬁlm a N-to-Al atomic ratio the same as the atomic arrival ratio of N-
to-Al in the depositing ﬂuxes. The arrival ratio of N-to-Al was converted from the i/a

ratio using a factor of 1.787 [Netterﬁeld, 1988].

Oxygen atoms from the impinging of the residual oxygen and water molecules in
the deposition chamber were incorporated into the ﬁlm, bonded in aluminum nitrate.
After the ﬁrst stage of deposition, the ﬁlms consisted of aluminum nitride, aluminum

nitrate, and metallic aluminum.

At the second stage of the ﬁlm deposition, all the metallic aluminum component
in the ﬁlm was oxidized by the impingement of the residual oxygen and water molecules,
in the absence of the aluminum and nitrogen ﬂuxes. Hence, after the second stage of
deposition, the ﬁlm consisted of three molecular components: aluminum nitride,

aluminum nitrate, and aluminum oxide.

The following two sections discussed in detail the composition of nitride ﬁlms
deposited at both 593 K and 793 K. Averaged data were used in estimations. All

components in the ﬁlms in both deposition stages were assumed stoichiometric.

Film Deposited at 593 K with an i/a Ratio of 0.56

The nitride ﬁlm deposited at 593 K had an averaged i/a ratio of 0.56 that gave a

N-to-Al atomic arrival ratio of unity [Netterﬁeld, 1988]. The ﬁlm composition was

92

expressed as Al [N] OJr after the ﬁrst stage of deposition. The oxygen content was

originated from the impingement of oxygen and water molecules in the chamber.

During deposition, the partial pressures of oxygen molecules (P02 ) and water
molecules (P1720) in the deposition chamber were both 1.3x10'S Pa. The averaged arrival
rate (1.) of the nitrogen ions was 4.85x10l4 (ions-cm’Z-sec'l). The ambient temperature
(T) was 300 K. The i/a ratio was 0.56. Using equations derived in Chapter 4, the oxygen

atomic concentration in the ﬁlm after the ﬁrst stage was calculated as:

1
1

+—
1.787-i/a

 

C0 = C32 + C320 = (5.21x10l9P02 + 3.27 x 10'9PHZO)

1.0 >77

-5
=(5.21x10'9 + 3.27 x10”) 1'3 x 10 = 0.07 = 7.0 at%.

4.85 x 10”(1 + 0.99)J300

Therefore, after the ﬁrst stage of deposition, the atomic composition of the ﬁlm was 46.5

at% aluminum, 46.5 at% nitrogen, and 7.0 at% oxygen.

After the ﬁrst stage of deposition, the ﬁlm was a mixture of aluminum nitride
(AlN), aluminum nitrate (A1(N03)3), and metallic aluminum (Al). For every 100 atoms in
the ﬁlm, 46.5 atoms were aluminum, 46.5 for nitrogen, and 7 for oxygen, respectively.
The 100 atoms formed a molecules of AlN, b molecules of A1(N03)3, and c atoms
(molecules) of Al. Hence, the ﬁlm components were: aAlN + bAl(N03)3 +ch1. The

amount of individual atoms in the ﬁlm was expressed as:

Aluminum Atom: a + b + c = 46.5
Nitrogen Atom: a + 3b = 46.5
Oxygen Atom: 9b = 7

Therefore, the number of molecules formed from every 100 atoms in the ﬁlm was

44.1 for AlN (a), 0.8 for A1(N03)3 (b), and 1.6 for Al (c) after the ﬁrst stage of deposition.

93

For every 100 atoms in the ﬁlm after the ﬁrst stage, 1.6 atoms were metallic
aluminum. At the second stage of ﬁlm deposition, the 1.6 atoms of metallic aluminum
component was oxidized into aluminum oxide (A120 3), incorporating an additional of 2.4
oxygen atoms into the ﬁlm. Within the 102.4 (=100+2.4) atoms, 46.5 of them were
aluminum, 46.5 for nitrogen and 9.4 for the oxygen. These 102.4 atoms in the ﬁlm after
the second stage would form 44.1 molecules of AlN, 0.8 molecules of A1(N03)3, and 0.8

molecules of A1203.

Therefore, the nitride ﬁlm deposited at 593 K with an averaged i/a ratio of 0.56
was estimated to consist of an atomic composition of 45.4 at% aluminum, 45.4 at%

nitrogen, and 9.2 at% oxygen. The molecular composition of the ﬁlm was estimated of

1.8% A1(N03)3, 1.8% A1203, and 96.4% AlN.

Film Deposited at 793 K with an i/a Ratio of 0.53

The nitride ﬁlm deposited at 793 K had an i/a arrival ratio of 0.53 that was
equivalent to an atomic N-to-Al ratio of 0.95 [Netterﬁeld, 1988]. After the ﬁrst stage of
deposition, the ﬁlm composition was expressed as AllN0.95Ox. The oxygen content was

originated from the impingement of the oxygen and water molecules in the chamber.

During the deposition, the partial pressures of oxygen molecules (P02 ) and water
molecules (P1720) in the deposition chamber were both 1.3x10'5 Pa. The averaged arrival
rate (1.) of nitrogen ions was 3.4x10'4 (ions-cm‘zosec'l). The i/a ratio was 0.53. The
ambient temperature (7) was 350 K. Using the equation derived in Chapter 4, the oxygen

atomic concentration in the ﬁlm after the ﬁrst stage was estimated as the following:

94

l

1,0 4N?

+
1.787 . i/a

 

C0 = C32 + C520 = (5.21><10'9P02 + 3.27 x10'9PH20)

-5
=(5.21x10l9 + 3.27 x10”) 1‘3 x 10 = 0,059 —_- 59 at%,

4.85 x10”(l+1.06)\[350

Thus, after the ﬁrst stage of deposition, the atomic composition of the ﬁlm was estimated

of 48.3 at% aluminum, 45.8 at% nitrogen, and 5.9 at% oxygen.

After the ﬁrst stage of ﬁlm deposition, the ﬁlm was a mixture of aluminum nitride
(AlN), aluminum nitrate (A1(N03)3), and metallic alurrrinum (Al). For every 100 atoms in
the ﬁlm, 48.3 atoms were aluminum, 45.8 for nitrogen, and 5.9 for oxygen, respectively.
The 100 atoms formed a molecules of AlN, b molecules of A1(N03)3, and c atoms
(molecules) of Al. Hence, the ﬁlm components were: aAlN + bAl(N03)3 + CA]. The

amount of individual atoms in the ﬁlm can be expressed as:

Aluminum Atom: a + b + c = 48.3
Nitrogen Atom: 0 + 3b = 45.8
Oxygen Atom: 9b = 5.9

Therefore, the number of molecules formed from every 100 atoms in the ﬁlm was

43.8 for AlN (a), 0.7 for A1(N03)3 (b), and 3.8 for Al (c) after the ﬁrst stage of deposition.

For every 100 atoms in the ﬁlm after the ﬁrst stage, 3.8 of them were metallic
aluminum. At the second stage of ﬁlm deposition, the 3.8 atoms of metallic aluminum
component in the ﬁlm was oxidized into aluminum oxide (A1203), incorporating an
additional of 5.7 oxygen atoms into ﬁlm. Within the 105.7 (=100+5.7) atoms after the
second stage of deposition, 48.3 of them were aluminum, 45.8 were nitrogen, and 11.6
were the oxygen. These atoms formed 43.8 molecules of AlN, 0.7 molecules of A1(N03)3,

and 1.9 molecules of A1203.

95

After deposition, therefore, the nitride ﬁlm deposited at 793 K with an averaged
i/a ratio of 0.53 was estimated to have an atomic composition of 45.7 at% aluminum,
43.3 at% nitrogen, and 11.0 at% oxygen. In terms of molecular concentration, the ﬁlm

consisted of 94.4% AlN, 1.5% A1(N03)3, and 4.1% A1203.

96

BIBLIOGRAPHY

97

BIBLIOGRAPHY -

Ballufﬁ, R.W., Blakely, J.M., Thin Solid Films, Vol. 25, pp. 363, 1975.
Behrisch, R., ed., Sputtering by Particle Bombardment 1., Springer-Verlag, 1981.
Behrisch, R., ed., Sputtering by Particle Bombardment 11., Springer-Verlag, 1983.

Behrisch, R., Wittmaack, K., ed., Sputtering by Particle Bombardment III., Springer-
Verlag, 1991.

Brindley, P. K., High-Temperature Ordered Intermetallic Alloys. II , pp. 419, 1987.

Broutman, LI. and Krock, R.H., Composite Materials Vol. 1: Interfaces in Metal Matrix
Composites, ed. by Metcalfe, A.G., Academic Press, 1974.

Bunshah, R.F., Deposition Technologies for Films and Coatings. Developments and
Applications, ed. by Bunshah, R.F., et. al., pp. 83, 1984.

Bunshah, R.F., Deshpandey, C., in Physics of Thin Films, Vol. 13, ed. by Francombe,
M.H., Vossen, J.L., Academic Press, New York, 1987.

Cai, Z., Thesis for the Degree of MS, Michigan State University, 1992.

Cairns, J.A., Nelson, R.S., Baunﬁeld, R.W., British Patent 55466/72, 1972.
Chambers, A., Fitch, R.K., Halliday, B.S., ed., Basic Vacuu_m Technology, 1989.
Chou, T.C., Joshi, A., Wadsworth, J., J. Vac. Sci. Technol, Vol. 9(3), pp. 1525, 1991.
Chou, T.W., Kelly, A., and Okura, A., Composites, Vol. 16, pp. 187, 1985.

Chou, T.C., Nieh, T.G., (a), Scripta, Metall, Vol. 25, pp. 2059, 1991.

Chou, T.C., Nieh, T.G., (b), High-Temp. Ordered Intermetallic Alloys IV, ed. by Johnson,
L.A., et.a1., pp. 1045, 1991.

Gardner, D.S., Flinn, P.A., IEEE Transa.on Electron Devices, Vol. 35(12), pp. 2160,
1988.

Gilmer, G.H., and Farrell, H.H., (a), J. Applied Phy., Vol. 47(9), pp. 3792, 1976.
Gilmer, G.H., and Farrell, H.H., (b), J. Applied Phy., Vol. 47(10), pp. 4373, 1976.

Glang, R., in Handbook of Thin Film Technology, ed. by Maissel, LL, and Glang, R.,
McGraw-Hill, New York, 1970.

98

Goshtagore, R. N ., J. Appl. Phys, 40, pp. 4374, 1969.
Grayson, M., ed., Encyclopedia of Composite Materials and Components, 1983.

Gupta, D., Camplell, D.R., H0, PS, in Thin Films - Interdiffusion and Reaction, ed. by
Poate, J.M., Tu, K.N., Mayer, J.W., 1978.

Gupta, D., H0, PS, ed., Diffusion Phenomena in Thin Films and Microelectronic
Materials, Noyes Publications, 1988.

Hall, E.L., Kouh, Y.M., Jackson, M.R., Mehan, R.L., Met. Trans. A., Vol. 14A, pp. 781,
1983.

Harper, J .M.E., Cuomo, J.J., Kaufman, H.R., Ann. Rev. Mater. Sci, Vol. 13, pp. 413,
1983.

Harper, J.M.E., Cuomo, J .J ., Kaufman, H.R., J. Vac. Sci. Technol, Vol. 21(3), pp. 737,
1982.

Harper, J .M.E., Cuomo, J.J., Hentzell, H.T.G., J. Appl. Phys, Vol. 58, pp. 550, 1985.
Hentzell, H.T.G., Harper, J .M.E., Cuomo, J.J., J. Appl. Phys, Vol. 58, pp. 556, 1985.
Hillegas, W., Solohill Eng, Ann Arbor, Michigan, Private Communication, 1990.

Holland, L., Science _a_nd Technology of Surface Coating, ed. by Chapman, B.N.,
Academic Press, 1974.

Jackson, M.R., Mehan, R.L., Davis, A.M., Hall, E.L., Metall. Trans. A., Vol. 14A, pp.
355, 1983.

Kattelus, H.P., Nicolet, M.-A., in Diffusion Phenomena in Thin Films Cid
Microelectronic Materisals, ed. by Gupta, D., and H0, PS, pp. 432, 1988.

Kaufman, H.R., Cuomo, J .J ., Harper, J.M.E., J. Vac. Sci. Technol, Vol. 21(3), pp. 725,
1982.

Kaufman, H.R., Robinson, R.S., Operation of Braod-Beam Sources, Commonwealth
Scientiﬁc Corporation, 1987.

Kingery, W.D., Bowen, H.K., Unlmann, D.R., ed., Introduction to Ceramics, second
edition, John Wiley & Sons, 1976.

LeClaire, A.D., Brit. J. Appl. Phys, Vol. 14, pp. 351, 1963.
Loretto, M.H., ed., Electron Beam Analysis of Materials, Chapman and Hall, 1984.

Lowell, C.E., Barrett, C.A., Whittemberger, J.D., Intermetallic Matrix Composites, ed. by
Anton, D.L., et. al., pp. 355, 1990.

99

 

Mattox, D., Materials Research Society Short Course on Films and Coating Deposition
Techniques, 1989.

Mehan, R.L., Bolon, R.B., J. Mat. Sci, Vol. 14, pp. 2471 - 2481, 1979.

Metcalfe, A.G., Composite Materials, Vol. 4: Metallic Matrix Composties, ed. by
Kreider, K.G.,.pp. 269,1974.

Moulder, J .F., Stickle, W.F., Sobol, P.E., Bomben, K.D., Handbook of X-ray
Photoelectron Spectroscopy, ed. by Chastain, Perkin-Elmer Corporation, 1992.

Movchan, B.A., Demchishin, A.V., Phys. Met. Metallogr., Vol. 28, pp. 83, 1969.

Netterﬁeld, R.P., Muller, K.-H., McKenzie, D.R., Goonan, M.J., Martin, P.J., J. Appl.
Phys, Vol. 63(3). Pp. 760, 1988.

Nicolet, M.-A., Thin Solid Films, Vol. 52, pp. 415, 1978.

Nieh, T.G., Stephens, J .J ., Wadsworth, J., Liu, C.T., Interfaces in Polymer, Ceramic, and
Metal Matrix Composites, ed. by Ishida, H., pp. 215, 1989.

Nourbakhsh, S., Margolin, H., and Liang, F. L., Metall. Trans. A, Vol. 20A , pp. 2159,
1989.

Ohring, M., ed., The Materials Science of Thin Films, 1991.

Poluboyarinov, D.N., Gordova, M.R., Kuznetsova, I.G., Bershadskaya, M.D., Avetikov,
V.G., Inorganic Materials, Vol. 15(11), pp. 1617, 1980.

Pope, D. P., High-Temperature Ordered Intermetallic Alloys II., pp. 3, 1987.
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Russ, J.C., ed., Fundamentals of Energy Dispersive X-ray Analysis, Butterworths & Co.
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Schoutens, J. E., Tempo, K., Introduction to Metal Matrix Copmposite Materials, DOD
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Smallman, R.E., Modern Physical Metallurgy, Butterworths, 1985.

100

Stelmack, L.A., Thurman, C.T., Thompson, G.R., Nuclear Instruments and Methods in
Physics Research, B, Vol. 37/38, pp. 787, 1989.

Taylor, J.A., Rabalais, J.W., J. Chem. Phys, Vol. 75, pp. 1735, 1981.

Thornton, J.A., Ann. Rev. Mat. Sci., Vol. 7, pp. 239, 1977.

Thornton, J .A., Hoffman, D.W., Thin Solid Films, Vol. 171, pp. 5, 1989.

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- Walters, R. P., Covino, B. 8., Jr., Metall. Trans. A., 19A, pp. 2163, 1988.
Westwood, W.D., Materials Research Society Bulletin, pp. 46, Dec. 1988.
Whipple, R.T.P., Philo. Mag, Vol. 45, pp. 1225, 1954.

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453, 1989.

Zangvil, A., Ruh, R., Mat. Sci. Eng, Vol. 71, pp. 159, 1985.

101

FIGURES

102

lg(D)

 

 

High Temperature Regime Low Temperature Regime
(Lattice Diffusion Domination) l (Grain-Boundary Diffusion Domination)

 

 

 

 

 

 

 

 

  

I‘\ Division between

Diffusion Dominations

 

 

 

 
  
 
 
  

 

 

 

D (combined)

 

\

D(grain-boundary),
with low Q

 

 

 

 

 

 

D(lattice),
with high Q

 

 

 

 

l/T

Figure 2.1. Illustration of Temperature Regimes of Dominant Diffusion
Mechanisms.

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

6 :T/Tm= 0.6 :T/rm: 0.5
4: b : b
s 2- : /
BB 0: -
3 -2: L d I L d
-4: I
-6l-l I l l l l l I l l l l -l I l l l l 1 l I l l
6:T/Tm= 0.4 :T/THF 0.3
4C I
c _ -
a 2- b - b
an 0: :
..1 -2_ d :
-4: L : d
-6 l L] l LLIJ l l 1 l I I l l l l l 1 LI I
0 2 4 6 8 1012 0 2 4 i 8 10 12
Logpd Logpd

 

 

 

 

 

 

Figure 2.2. Regimes of Dominant Diffusion Mechanism in FCC Metal
Films [after Ballufﬁ, 1975].
1: Grain Size (in cm) pd: Dislocation Density (in cm‘z)
L: Lattice Diffusion Dominant b: Grain-Boundary Diffusion Dominant
d: Dislocation Diffusion Dominant

103

 

l E+0
— Film Model

- - - - Bulk Model

 

 

 

 

1E-2 . . . .

(y/Yo)6/5

Figure 2.3. A Comparison of Concentrational Proﬁles between the Film
and Bulk Models [after Gilmer, 1976a].

Yo: Film thickness y: Distance from the source interface

The high diffusant concentration in thin ﬁlm is attributed to the boundary
condition of a free surface imposed to the thin ﬁlm model.

104

src M643, A1181; N15_4A11 312 MA] Ni 3A1

ma 7. / /

 

j
f

\\
Si

--u-t

E!

j.
. $11)

I
I
I
I
I
I
I
I
I
I

y.

 

 

 

 

The Approximate Position
of the Original Interface

 

IOO-ISOum

Figure 2.4. Schematic Presentation of Reaction Zone between SiC and
y-Ni3Al at 1273 K for 23 Hours [Chou, 1991b].

SiC Ni14 A198i2 NiAl
+C(graphite)

 

r~
Q

 

 

 

 

T
.1.

14m

Figure 2.5. Schematic Presentation of Reaction Zone between SiC and
B-NiAl at 1573 K for 20 Hours [Chou, 1991a].

105

 

 

 

 

 

273 423 573 723 873
0.5 ;
.§ 0.4 """""" ‘- """""
5 :
5‘ 0 3r """"" 1 """""""""
:3 I
E 0.2 ---------- 4. --------
...I r
S l
0.1 ---------- . - - -
0 ' l ' 600
0 150 300 450

Figure 2.6. A Comparison of Linear Thermal Expansion of AlN, SiC, and
A1203 [after Carborumdum Co., 1992].

 

     

 
    
 

 

 

 

 

 

T (K)
273 373 473 573
200 T ;
9 IN I I
E ' J
g 150 """"""" : """""" . """""""
jg. : .
E 1001 ------- , '''' :- """""""" """""""
g SIC I r
E ' I
5 50 .............. - .......... 3 --------------
.3 A1203 : : _
E g : :
'ﬁ 0 . -
0 100 200 300

T (°C)

Figure 2.7. A Comparison of Thermal Conductivity of AlN, SiC, and
A1203 [after Carborumdum Co., 1992].

106

 

Substrate
/

To Vacuum Pump
Ion
Source \ Electron Beam
Inlet of Inert or

Reactive Gas / Deposrtrng Material

—-'9I=

 

 

—— E-beam Source

\ Water-cooled Hearth

 

 

Figure 2.8. Schematic Conﬁguration for Ion-Beam-Assisted Electron-
Beam Evaporation.

P .
. ;Q\ “I
I2: " a

.. ~ ' .4 1'!

m °
111111

e . I‘Q'Sﬁlﬁe ‘0

- '0 3
.IA.

 

0.1 0.2 0.3 0.4 0.5 0.6 0.7

Figure 2.9. Schema of Zone Structures in Evaporated Thin Films [after
Movchan, 1969].
Films with Zone 1 structure consist of small tapered crystals with voided
boundaries. Zone T structure has a mixture of small and large grains,
which produces a dense ﬁlm. Zone 11 structure offers a columnar grains
and dense grain boundaries. Zone 111 structure is only formed at high
deposition temperature, consisting of large equiaxed grains.

107

U U in

 

Figure 3.1. A SEM Photo of an As-Received SiC Particle.
The SiC particle was embedded in an Indium foil. The
nominal SiC particle size was 60 11m.

108

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110

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Intermetallic L _ Al 1 t
alloy plate 203 P a e
E —— Diffusion couple
SiC

particles I::::::l— 23203131,...

Intermetallic ll

Loading

 

 

alloy plate
Loading Loading
(a) Diffusion Couple (b) Testing Conﬁguration

Figure 3.4. Conﬁguration for Hot-Pressing.

111

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XPS Spectrum

Curve-ﬁt

 

 

534 533' 5'32f5'31‘5'30'5'29'5'28 527 526
Binding Energy, eV

(a). Oxygen Peak

 

 
    
  

XPS Spectrum

Curve-ﬁt

 

 

      

 

l A

'78 77 7‘6-7.5 74 73 ‘72 71
Binding Energy, eV

(b). Aluminum Peak

Figure 4.7. XPS Peaks from an Aluminum Oxide Film Deposited at
Room Temperature, after Curve-ﬁt.

Both peaks were generated from aluminum oxide. The composition was
calculated to be Al33062. The sample was determined to be nominally
stoichiometric aluminum oxide.

118

 

6O

50 +

.h
C
l
I

Atomic Percentage
b)
o
l
T

 

 

 

20 ~~
10 -~
0
0 l l + + l l
O 10 20 3O 4O 50 6O 7O

Sputtered Depth from Surface (nm)

Figure 4.8. XPS Concentrational Proﬁle of an Aluminum Nitride Film
Deposited at 793 K with an i/a of 0.42.
The N-to-Al ratio in the ﬁlm were corresponding to the 170 ratio of 0.42.
A 5 — 7 at% oxygen content existed throughout the ﬁlm. The higher
oxygen content near ﬁlm surface region was attributed to the oxidization
after exposing the ﬁlm to air.

119

 

 

 

XPS Spectrum

’I 0‘ Curve-ﬁt

" (Nitrate Peak at

532.7eV)
, \
’ l
l M, \M
A K : t c : - c : “ h.
36 535 534 533 532 531 530 529

Binding Energy, eV

1

J
v

 

(a). Oxygen Peak, with Curve-ﬁt

 

 

 

 

 

XPS Spectrum M
Curve-ﬁt [
\ Deconvoluted
’ \ Peak 1
// (Mixed Compounds
4 Peak at 73.8 eV)
Deconvoluted
I Peak 2
/ ( (Meztaélicv? Peak
l“ at 7 . c
» >-
76 75 ' 74 ' .43 T 72 71
Binding Energy, eV

(b). Aluminum Peak, after Deconvolution

Figure 4.9. XPS Peaks from an Aluminum Nitride Film Deposited at 793 K

with an i/a of 0.42, 40 nm below the Original Film Surface.

The XPS peaks of oxygen and aluminum indicated the existence of aluminum
nitrate and metallic aluminum components inside the nitride ﬁlm.

120

 

 

  
 
   
 
 
 
  

Curve-ﬁt Deconvoluted

Peak 1
(Nitrate Peak at

Deconvoluted
Peak 2

(Aluminum Oxide
Peak at 531.8 eV)

   
 
 
   
 
 
 
 

 

 

 

Binding Energy, eV

(3). Oxygen Peak, after Deconvolution

 

Curve-ﬁt —————--—""""
Deconvoluted

Peak 1
Mixed Compound
Peak at 74.1 eV)

Deconvoluted

Peak 2
(Metallic Al Peak
t 72.9 eV)

  

 

 

77 76 75 74 73 72 71 70
Binding Energy, eV

(b). Aluminum Peak, after Deconvolution

Figure 4.10. XPS Peaks from an Aluminum Nitride Film Deposited at

793 K with an i/a of 0.42, at Surface Region.

The XPS peaks of oxygen and aluminum indicated the existence of
aluminum nitrate, aluminum oxide, and metallic aluminum
components on the surface of the nitride ﬁlm.

121

 

 

XPS Spectrum

 
 
  
 
   
   
   

Deconvoluted

Peak 1
(Nitrate Peak at
533.3 eV)

Curve-ﬁt

Deconvoluted

Peak 2
(Aluminum Oxide
Peak at 53 l .7 eV)

 

 

 

540 538 536 534 532 530 528 526
Binding Energy, eV

Figure 4.11. The XPS Oxygen Peak from an Aluminum Nitride Film

Deposited at 593 K with an i/a of 0.60, Surface Region.

The oxygen XPS peak indicated the existence of aluminum nitrate and
aluminum oxide components on the surface of the nitride ﬁlm.

 

XPS Spectrum N

    
 

Curve-ﬁt
(Nitrate Peak
at 532.5 eV)

 

 

A.

540 538 536' 534' 532' 530 528 526
Binding Energy, eV

 

Figure 4.12. The XPS Oxygen Peak from an Aluminum Nitride Film
Deposited at 593 K with an 170 of 0.60, 20 nm below the

Original Film Surface.

The oxygen XPS peak indicated the existence of aluminum nitrate
component inside the nitride ﬁlm.

122

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123

7 F 9. 1‘1 9

 

Figure 4.14. Aluminum Nitride Film Deposited on SiC Particle at 593 K,
after Quenching in Liquid Nitrogen (77K) from 1273 K.
The nitride ﬁlm retained the attachment on a fractured SiC particle after
the annealing. The ﬁlm structure was globular, similar to ﬁlm’s as-
deposited structure.

124

 

Figure 4.15. Aluminum Oxide Film on a SiC Particle, after Air-cooling

from 1273 K to Room Temperature.
Grain grth occurred in the ﬁlm during the annealing, resulting in

micro-cracks in the ﬁlm.

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128

Figure 4.18 (Cont.)

 

(c). A SEM Image of the Film Surface with High Magniﬁcation

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C(Ni)/C(Al)

 

 

  

    

 
 
    

 

 

 

 

 

 

  

 

 

 

 

 

   

 

 

 

 

 

3.5
Substrate Nitride Film E
3 *- 3
l
2.5 «~
Film Surface
2 4 _
Distance
1.5 W =215 nm
i
1 w I
5
Distance i
0.5 w =225 nm Distance ;
A :25 um i
0 f 5 f i E Y Y , l—
900 800 700 600 500 400 300 200 100 O
(Distance)m ("my5 )
Figure 4.22. XPS Proﬁle of Ni Diffusing in an Aluminum Nitride Film

from a y-Ni3Al Source after a 1373 K130 Minutes Annealing.

The ﬁlm was deposited on a y-Ni3Al plate. The total distance of Ni
diffusing in the ﬁlm was about 200 nm (0.2 um).

133

 

 

(a). The Reaction Zone between SiC and MA]

 

Al

 

 

Si Ni
- M

l l
O 2 4 6 8
X-ray Energy, KeV

 

 

10

(b). EDX Spectrum from the Reaction Zone

Figure 4.23. Reaction Zone in a SiC Particle Reinforced B-NiAl Matrix

Composite with Uncoated SiC Particles.
The reaction zone was found in between NiAl matrix and the remnant of

the SiC particles.

134

Reaction Zone

 

Figure 4.24. Reaction Zone in a SiC Particle Reinforced y-Ni3Al Matrix
Composite with Non-coated SiC Particles.
(Magniﬁcation: 500x)
The reaction zone was found at the location previously occupied by the SiC
particles. No remnant of the SiC particles was found in the composite.

135

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136

Figure 4.25 (Cont.)

 

Si

 

 

 

 

EJLW IN:

0 2 4 6 8 10
X-ray Energy: KeV

 

(c). EDX Spectrum from the Revealed SiC Surface Shown in Figure 4.25 (b)

137

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138

Figure 4.26 (Cont.)

 

Ni
Si

 

 

 

 

 

 

A) l L/\

I l_
l I
4 6 8

 

 

 

N‘-
-.

0

X-ray Energy: KeV

(c). EDX Spectrum from the Reaction Products Revealed at the
Film Damaged Area

139

 

 

 

Figure 4.27. Reaction Zone in a SiC Particle Reinforced y-Ni3Al
Matrix Composite with Aluminum Nitride Film
Deposited at 593 K as a Diffusion Barrier.
(Magniﬁcation: 500x)
The reaction zone was found at the location previously occupied by
the SiC particles. No remnant of the SiC particles was found in the
composite.

140

 

Figure 4.28. Reaction Zone in a SiC Particle Reinforced B-NiAl Matrix
Composite with Aluminum Nitride Film Deposited at 793 K
as a Diffusion Barrier.

The reaction zone was found in between NiAl matrix and the remnant of
the SiC particles. No evidence of stalling the reaction by the imposition
of the nitride barn'er ﬁlm was found.

141

 

Figure 4.29. Reaction Zone in a SiC Particle Reinforced 'Y-Ni3Al Matrix
Composite with Aluminum Nitride Film Deposited at 793 K
as a Diffusion Barrier.
The SEM sample was prepared by fracturing the composite sample in tension.
The reaction zone was found at the location previously occupied by the SiC
particles. No remnant of the SiC particles was found. No evidence of stalling
the reaction by the imposition of the nitride barrier ﬁlm was found in the
composite.

142

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