Low Energy
j Qjrowin ^
1.1
A R Gonzalez-Elipe F Yubero J M Sam
Imperial College Press
Low
Energy
Ion Assisted ilm Growth
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Low
Energy
Ion Assisted Film Growth
A 1 Sisiiilei-llipe F Tuber© ImtHuto de Cienca d$ Materiales de Semlla (CSIC-M Sevilla), Spain
J in SUM UniversidodMonoma de Madrid, Spsln
Imperial College Press
Published by Imperial College Press 57 Shelton Street Covent Garden London WC2H 9HE Distributed by World Scientific Publishing Co. Pte. Ltd. 5 Toh Tuck Link, Singapore 596224 USA office: Suite 202, 1060 Main Street, River Edge, NJ 07661 UK office: 57 Shelton Street, Covent Garden, London WC2H 9HE
British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library.
LOW ENERGY ION ASSISTED FILM GROWTH Copyright © 2003 by Imperial College Press All rights reserved. This book, or parts thereof, may not be reproduced in any form or by any means, electronic or mechanical, including photocopying, recording or any information storage and retrieval system now known or to be invented, without written permission from the Publisher.
For photocopying of material in this volume, please pay a copying fee through the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, USA. In this case permission to photocopy is not required from the publisher.
ISBN 1-86094-351-9
Printed by Fulsland Offset Printing (S) Pte Ltd, Singapore
Contents
Foreword
xiii
CHAPTER 1: BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS WITH SOLID TARGETS
1
1.1. Introduction
2
1.2. Interatomic interaction
3
1.2.1. Atoms in condensed matter 7.2.2. Interaction of energetic ions with condensed matter: Interatomic potential 1.2.3. Power law approximations to the interatomic potential 1.3. Basic concepts in classical dynamics of binary elastic collisions 1.3.1. Ion energy loss rate 1.3.1.1. Nuclear stopping 1.3.1.2. Electronic stopping
3 6 10 12 14 16 19
1.4. Range of energetic ions in solids
20
1.5. Spatial distribution of deposited energy
25
1.6. Damage induced by ion bombardment
27
1.6.1. 1.6.2. 1.6.3. 1.6.4.
Primary knock-on atoms formation Spikes Thermal spikes Density of the deposited energy
1.7. Sputtering 1.7.1. Sputtering yield 1.7.2. Angular distribution of sputtered atoms 1.7.3. Energy distribution of the sputtered atoms v
27 30 32 33 35 36 38 39
VI
CONTENTS
1.8. Experimental parameters in IAD thin film growth 1.8.1. The ion to atom arrival ratio and the normalized energy concept 1.8.2. Ion momentum transfer
40 41 , 44
References
45
CHAPTER 2: ION ASSISTED METHODS OF PREPARATION OF THIN FILMS
47
2.1. Assistance of film growth with independent ion sources 2.1.1. Evaporation and ion bombardment of the growing film 2.1.2. Laser ablation and ion bombardment of the growing film 2.1.3. Dual ion beam deposition of thin films (DIBS) 2.1.4. Ion beam induced chemical vapour deposition (IBICVD)
47 49 52 55 58
2.2. Ion assisted deposition of thin films without independent ion sources 2.2.1. Ion plating 2.2.2. Ionised magnetron sputtering (IMS) 2.2.3. Filtered vacuum arc deposition (FVAD) 2.2.4. Ionised cluster beam (ICB) 2.2.5. Mass selected ion beam deposition (MSIBD)
60 60 64 66 69 72
2.3. Plasma immersion ion implantation 2.3.1. Plasma immersion ion implantation (PHI) 2.3.2. Plasma immersion ion deposition (PHD)
75 76 78
2.4. Broad beam ion sources 2.4.1. Kaufmann type ion sources 2.4.2. End-Hall ion sources 2.4.3. Filament-less ion sources
80 81 83 85
References
87
CONTENTS
vii
CHAPTER 3: EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
90
3.1. Ion beam effects during film growth
90
3.2. Nucleation and growth of thin films under ion bombardment
92
3.2.1.
Nucleation and growth of physical vapour deposited (PVD) thin films 3.2.2. Effects of ion bombardment on nucleation 3.2.3. Monitoring the surface defects and nucleation processes induced by ion bombardment 3.2.4. Description of Nucleation and Growth processes by analysis of STM/AFM images 3.3. Topography and surface and interface roughness 3.3.1. Grain size 3.3.2. Surface roughness 3.3.3. Step and surface coverage 3.3.4. Surface roughness of thin films grown by IBD 3.3.5. Interface roughness 3.3.6. Monitoring the interface roughness by X-ray reflectometry 3.3.7. Epitaxial growth of thin films 3.4. Interface mixing 3.4.1. Mixing in thick films and bulk materials induced by high energy ions 3.4.2. Interface mixing in IAD thin films 3.4.3. Monitoring interface mixing by TEM/EELS
92 93 96 97 99 99 100 103 103 104 106 108 108 109 110 Ill
3.5. Densification of thin films 3.5.1. Columnar growth in PVD thin films 3.5.2. Densification in IAD thin films 3.5.3. Evolution of density and crystallinity with ion energy and I/A ratio
113 113 114 116
3.6. Defect generation 3.6.1. Formation of defects in IAD thin films 3.6.2. Surface and bulk defects as a function of beam energy
119 119 120
CONTENTS
Vlll
3.6.3. Defects and control of the microstructure of thin films by annealing treatments 3.6.4. Inert gas incorporation
122 123
3.7. Amorphisation, crystallinity and phase transformations 3.7.1. Amorphisation in IAD thin films 3.7.2. Effect of temperature on crystallisation 3.7.3. Amorphisation and phase transformation phenomena. Stabilisation of unstable phases 3.7.4. Monitoring the degree of amorphisation in IAD thin films
128 128
3.8. Compound formation by IAD 3.8.1. Control of stoichiometry in IAD thin 3.8.2. Metastable phases of nitride thin films
130 131 133
films
3.9. Texture development 3.9.1. Monitoring the texture in IAD thin films by XRD: Basic definitions 3.9.2. Texture inPVD thin films 3.9.3. Texture evolution in IAD thin films and process parameters 3.9.4. Models for texture development 3.9.5. Biaxial orientation 3.9.6. Applications of textured thin films
125 126 126
137 137 142 143 147 150 153
3.10. Influence of ion assistance on thin film stress 3.10.1. Basic concepts on stress 3.10.2. Distribution of stress between substrate and thin film 3.10.3. Thermal stress in thin films 3.10.4. Intrinsic stress in PVD thin films 3.10.5. The stress in IAD thin films: Dependence on experimental parameters 3.10.6. Compressive stress in IAD thin films 3.10.7. The stress in IAD thin films: Models
154 154 154 156 156
3.11. Improvement of adhesion in IAD thin
165
References
films
158 161 162
168
CONTENTS
ix
CHAPTER 4: APPLICATIONS OF IAD PROCESSING
173
4.1. Tribological coatings 4.1.1. Hard and wear resistant coatings 4.1.2. Solid lubricant coatings
174 175 177
4.2. Corrosion resistant coatings 4.2.1. Metal coatings 4.2.2. Oxide and nitride coatings 4.2.3. Corrosion protection of magnesium alloys 4.2.4. Zinc and zinc alloys
179 181 182 183 184
4.3. Modification of biomaterials 4.3.1. Fretting wear and damage 4.3.2. Corrosion protective coatings 4.3.3. Hydroxiapatite 4.3.4. Biocompatibility
185 185 186 186 187
4.4. Metallisation of polymers
188
4.5. Optical coatings 4.5.1. Dielectric oxide films 4.5.2. Fluoride thin films 4.5.3. Narrow bandfdters 4.5.4. Rugate 4.5.5. Transparent conducting
films
190 191 196 197 199 201
films
204 205 206 208 209
4.6. Magnetic thin films 4.6.1. Thin metallic 4.6.2. Magnetoresistive materials 4.6.3. Reading/writing magnetic heads 4.6.4. Hard bias magnetic thin References
filters
films
211
X
CONTENTS
CHAPTER 5: DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
216
5.1. Diamond-like carbon
217
5.2. Characterization methods and related properties 5.2.7. Hydrogen concentration 5.2.2. Atomic structure (electron and neutron scattering) 5.2.3. sp3/sp2 bonding fraction (Raman, NMR, EELS/XAS) 5.2.4. Density 5.2.5. Cross sectional structure and in-depth composition (TEM.EELS)
219 220 220 221 224
5.3. DLC deposition methods
228
5.4. Influence of the deposition parameters on the sp3 bonding fraction and related properties 5.4.1. Influence of the ion energy 5.4.2. Influence of the substrate temperature 5.4.3. Influence of other deposition parameter
231 231 234 238
5.5. Stress in DLC
films
239
5.6. Properties and applications of the DLC films 5.6.1. Mechanical and tribological properties 5.6.2. Optical and electronic properties
240 242 243
5.7. Cubic Boron nitride films
245
5.8. Characterization of c-BN 5.5.7. Stoichiometry 5.8.2. XRD diffraction 5.8.3. Phase identification by FTIR spectroscopy and EELS/XAS 5.8.4. Microstructure by TEM
245 246 246 247 250
5.9. c-BN deposition methods
251
5.10. Influence of the deposition parameters
252
227
CONTENTS
xi
5.11. Stress
255
5.12. Properties and applications of c-BN films 5.12.1. Tribological properties 5.12.2. Optical and electrical properties
257 257 258
5.13. Modelling the growth of sp3 bonded materials (ta-C, ta-C:H and c-BN) 5.13.1. The preferential sputtering model 5.13.2. The stress models 5.13.3. Models involving a thermal spike mechanism 5.13.4. Subplantation models
258 260 260 261 262
5.14. Related materials (CNX, B-C-N)
264
References
268
Acronyms List
275
Subject Index
279
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Foreword
This book deals with the use of accelerated ion beams to assist the growth of thin films. It was recognised as early as the seventies that the bombardment of a growing film with a beam of accelerated ions induces significant changes in its properties and that these changes have many potential applications. The term "Ion beam assisted deposition" (IBAD) and the more general one "Ion assisted deposition" (IAD) were coined to characterise all the procedures of preparation of thin films that, in one way or another, use this type of approach. Since then, the modalities of this technology and its applications, in very different scientific and industrial areas, have expanded to constitute what is now a mature technique, useful for a tailored synthesis of thin films. Within a general perspective, several books and reviews have dealt with the phenomena involved during the IAD of thin films. Itoh's book of 1989 and some review papers by Smidt (1990), Hirvonen (1991) and Ensinger (1994, 1995, 1997) are noted examples of such publications. This literature covers the advances made during the seventies and eighties in the development of IAD techniques. Much research effort has contributed to new developments and to a more thorough understanding of the basic phenomena involved during IAD of thin films, not only providing an empirical perspective on the effects of ion bombardment on thin film properties, but permitting a clear interpretation of these phenomena on an atomistic scale. The present book aims to provide a comprehensive description of the basic phenomena involved in IAD processes, the different techniques of preparation of thin films that can be considered as ion assisted methods, and some of the applications of the prepared thin films in different fields of science and technology. Throughout this presentation, emphasis is put on results which appeared in scientific literature during the last decade, given that the most important contributions made before are properly discussed in these previous publications. This book is written with the intention that it serve as an introductory manual for researchers, post-graduates and engineers from industry with little or no experience in IAD thin films. Nonetheless, some basic principles of the interaction xm
XIV
FOREWORD
of ions with solid targets are also reviewed with the intention of providing a good scientific basis that will permit a sound and justified presentation of results and phenomena. In this respect, it is hoped that the book will also serve to provide researchers active in this scientific area with a general scheme and practical ideas for the improvement of their investigation in this interesting field of material science and technology. The book is divided into five chapters dealing with, respectively, the basic principles of the interaction of accelerated ions with matter; a description of the different techniques relying on the IAD concept; the changes experienced by the thin films when subjected to ion bombardment; some applications of the IAD thin films in different fields of science and technology and a description of the major issues related with two sets of materials for whose synthesis the use of IAD procedures is essential (i.e., c-BN and diamond and related materials). Whenever possible, results by updated methods of characterisation of thin films are included as examples of the possibilities of the IAD procedures. In some cases, this presentation of experimental results is accompanied by a brief presentation of the technique, so that non-experts will also be able to understand the main message of the proposed case. We hope that the reader will find this manuscript interesting and, what would be even more satisfactory for the authors, useful for their own investigations. Finally, we would like to thank Jose A. Rodriguez for his good job preparing the drawings of the different deposition methods included in Chapter 2, and the support provided by the author's research institutions (i.e., CSIC, Universidad de Sevilla and Universidad Autonoma of Madrid).
Seville and Madrid May 2002
CHAPTER 1 BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS WITH SOLID TARGETS
The assistance of the deposition of thin films with ion bombardment produces significant modifications in their characteristics and therefore in their final properties. Thus, densification, adhesion to a substrate, intrinsic stress, grain size, texture (preferential orientation of certain crystallographic planes), or even changes in the growing phases can be affected and controlled by ion-assisted deposition (IAD) of thin films. To understand how ion bombardment induces such effects, the physical basis of the interaction of energetic particles with condensed matter needs to be established. By energetic particles we mean those with kinetic energies typical of IAD processes, i.e., from a few tens to a few thousands of eV. Mostly, these particles are ions or at least particles whose kinetic energy is supplied when they are charged, before their interaction with the growing film. Throughout this book we use the term "ion" for the impinging particles that assist the film growth, but we should bear in mind that energetic neutral particles would show similar effects to the corresponding ions. As an energetic ion traverses a solid, it interacts with the electrons and the nuclei of the atoms forming its structure. This interaction results in the deflection of the ion trajectory from its original direction until it stops. The key point is that part of the initial kinetic energy of the ion is shared with the atoms of the solid due to collision events. Ion beam modification of materials can be explained by describing the ion-target atom interactions. Thus, through a proper description of such an interaction, it is possible to estimate the depth reached by the ions or the energy deposited along the ion track in the solid. In this chapter, the basic concepts of the interaction of energetic ions with condensed matter are introduced. The interaction between the energetic ions and target atoms is described through interatomic potentials. Concepts such as the stopping power, range of the penetrating ions, or damage energy are introduced and evaluated within simplified models. In addition, the consequences of the energy transfer of the impinging particles to the solid are also considered. Thus, the 1
2
Low ENERGY ION ASSISTED FILM GROWTH
formation of ion cascades, thermal spikes and the appearance of sputtering will be described. In all cases, we will try to give analytical expressions for the evaluation of the different magnitudes presented. These expressions do not pretend to be exact, but rather to give an indication of the dependencies on the energy and other characteristics of the ions (e.g. atomic/molecular weight, charge) and similar features for the rest atoms in the solid. Readers interested in a more exhaustive description of the interactions are referred to the works Lindhard et al. (1963), Winterbon et al. (1970), Sigmund (1981), Eckstein (1991) or Nastasi et al. (1996).
1.1. Introduction In this chapter, we introduce several magnitudes that enable an analytical description of the interaction between energetic ions and substrate atoms regarding mass transport and energy deposition. The basis of these interactions will be described by the interatomic potential between two particles. The effect of thermal vibrations, the concepts of the binding energy of a set of atoms and displacement energy will be introduced. Other concepts closely related to the description of the interatomic potential, such as screening function and screening lengths, will also be presented. To get a simplified view of the interaction, the so-called power law approach will be considered for the description of the interatomic potential and used extensively throughout this chapter. Binary collisions between impinging ions and rest atoms are always considered. These collisions will be characterised energetically by the so-called reduced energy of each particular collision. The use of this magnitude is convenient to sort out the type of interaction taking place, which in principle depends not only on the actual kinetic energy of the impinging ions, but also on the particular characteristics of the colliding particles. Ion energy losses can be due to "elastic" collisions with the rest atoms of the substrate and to the electronic excitations. These two effects will be quantified by using scattering probabilities or scattering cross sections for each type of interaction.
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
3
Once the physics of the interactions is established, we will proceed to describe the ion penetration range and damage energy distribution. These two magnitudes will give us a fairly clear idea of the size of the altered layer due to assistance with energetic ions and the expected effects in this altered layer due to the amount of energy deposited. As we will justify, atom relocation is a natural result of the deposited energy, and the damage introduced in the surface will be described by the formation of "spikes" or local atom rearrangements along the ion tracks. As a consequence of the deposition of energy in the altered layer, removal of material from the target (i.e., surface sputtering) takes place. The sputtering phenomena will be discussed and correlated with the energetic of the bombarding ion - target atom interactions. Finally, at the end of this chapter we will also present a description of the parameters that can be controlled in practice during IAD growth of thin films and their correlation with proposed models that justify experimental results obtained by the assisted growth. Throughout this book, we use the units typically considered by the thin film community. Thus, we will express lengths and thicknesses in nanometers (nm), kinetic energies of the particles or deposited energies in electron volts (eV). The mass and charge of the particles involved will be expressed in atomic units.
1.2. Interatomic interaction 1.2.1. Atoms in condensed matter Atoms in solids are located at their equilibrium positions, distributed according to the constraints of interatomic potentials V(r). In a solid at equilibrium, the interatomic potential must have a minimum at an equilibrium distance r0. The actual location of the atoms in a solid is a compromise between the attractive and repulsive forces acting on them. There are several models that have been proposed for the description of the interatomic potential of atoms in solids. Among them, the hardsphere potential, the square-well potential and the more realistic one proposed by Morse (1929) and Lennard and Jones (1924) can be mentioned. The latter can be expressed analytically in the form
4
Low ENERGY ION ASSISTED FILM GROWTH
V(r) =
(1.1) \
r
J
where ch c2, p and q (p > q) are constants. Figure 1.1 shows a schematic representation of the Lennard-Jones interatomic potential V(r). It is composed of positive and negative contributions that represent repulsive and attractive energies, respectively. The repulsive part vanishes more rapidly than the attractive, so that the addition of both contributions results in a minimum for the potential at the equilibrium position r0. Thus, according to this potential, if the distance between the two atoms is smaller than r0 the atoms will experience repulsion, while if r>r0 the atoms will experience attraction.
V(r) ', repulsive energy<* I ^
0
;
( r ^
atractive energy oc -
\
r
J
Figure 1.1. Lennard-Jones type interatomic potential V(r).
The depth of the potential well at r0 (i.e., £fc) is related to the binding energy of a solid Eb. This binding energy is proportional to the coordination number nc and to the minimum potential energy in the interatomic potential £ i; so that Eh « n„Eh
(1.2)
An implication of Eq. (1.2) is that the binding energy of an atom at the surface will be smaller than that of the same atom in the interior of a solid. This fact stems from a simple consideration of the lower coordination number of the atoms at
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
5
the surface in respect to that in the bulk. Then, it is possible to distinguish between surface binding energy Ebs and lattice binding energy Ebl. It is worth mentioning that Ebs is closely approximated by the heat of sublimation AHS. In practice, Eb and Eb take values between 2 and 8 eV depending on the type material (atoms involved, crystallographic structure, etc.) Another aspect to be considered is the thermal energy of the atoms in solids. It is given by KBT, where KB is the Boltzmann constant and T the temperature. At room temperature, it accounts for 0.024 eV, i.e., much less than the kinetic energy of the impinging ions considered in thin film IAD growth. Note that this kinetic energy may vary from a few tens to several thousands of eV depending on each particular case. Thus, we can advance that energy transfer by ion bombardment is locally much more energetic than thermal heating. The atoms in solids oscillate around their equilibrium positions due to their thermal energy. The shape of the interatomic potential well limits the amplitude of the vibrations as it is schematically represented in Figure 1.2. The length of the horizontal lines represents the amplitude of the vibrations, so that it increases as the temperature T, increases. Another consequence of thermal heating is that the mean distance between atoms will increase as the temperature increases. The shape and the depth of the potential well will be related to the thermal and elastic properties of
V(r) 0
I
mean atom position r
i r2 r<
I
^-—~ Ti
1 M l
r
\ W/^ \
W
f /
T
2
T
l
Figure 1.2. Effect of temperature 7", in the mean interatomic distance r, and amplitude of the thermal vibrations (represented as horizontal lines) of atoms in condensed matter.
6
Low ENERGY ION ASSISTED FILM GROWTH
the solid. Thus, for example, the increase of interatomic distance with increasing temperature (i.e., the thermal expansion) will be higher for solids with shallow asymmetric potential wells than with deep symmetric ones. Another concept related to the binding energy of an atom in a solid is the displacement energy Ed. It is defined as the energy required to displace an atom from its lattice site. Atom displacement usually results in defect formation as, for example, a vacancy-interstitial defect (Frenkel pair). The crystallographic structure of the solid imposes that the displacement barrier of lattice atoms is not uniform in all directions. It is lower in open directions (i.e., directions with high Miller indexes) than in directions where the atoms are more densely packed (i.e., directions with low Miller indexes). The displacement energies Ej vary from -5 eV for fragile materials as polymers to -25 eV for the case of metals. However these values may vary by a factor of 2 depending on the direction in which the displacement takes place (Nastasi et al., 1996). A phenomenological finding is that the Ed threshold is around 4-5 times the heat of sublimation AHS (T.E. Mitchell, 1976).
1.2.2. Interaction of energetic ions with condensed matter: Interatomic potential It is important to define the limits for the kinetic energy that will be considered in the following for the ion beams used in IAD. As mentioned before, thin film growth by IAD methods is usually assisted with ion beams from a few tens to several thousands eV. Therefore, we will consider accelerated ion beams with kinetic energies in the range 20 - 10000 eV. It is interesting to distinguish between bombardment with monoatomic ions and polyatomic molecular ions. The former is well represented by inert gas bombardment (i.e., Ne+, Ar+, or Xe+), where the kinetic energy of the ion can be identified with the kinetic energy of the particle interacting with the solid. However, in the latter case, the kinetic energy of the ion is shared among all the individual atoms in the molecule. This is relevant because as soon as one of these polyatomic ions contacts the surface of a solid, it splits into its individual atomic components. Thus, for example, in the usual case of oxygen or nitrogen (i.e., O^ or N2+) bombardment, the kinetic energy of each component atom can be considered as half of the total kinetic energy of the molecular ion.
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
7
The range of velocities of the considered ions is much smaller than the Bohr velocity of the electrons in their orbit, v0 - 2.2xl0 6 m/s. The velocity v of the ions expressed in m/s can be obtained from their kinetic energy E given in eV according to v = 13.9x10s
—
(1.3)
where M, is the molecular mass of the ion in atomic units. As a typical example, the corresponding velocity of Ar + the ions of 1 keV is ~7xl0 4 m/s, which is far too low to be considered relativistic (the speed of light is 3xl0 8 m/s). Therefore it is a good approach to imagine that the impinging ions in IAD processes "see" the solid as a set of fixed positive charges inserted in a sea of negative charge. To study the interaction between ions and a solid surface, we have to consider first the simplified idea of the collision of a moving particle with an atom at rest. To describe this interaction, the interatomic potential has to be identified. In general, it is assumed that a single interatomic potential function describes the entire collision process. In section 1.2.1, several model interatomic potentials are introduced. The equilibrium position of atoms in solids is described as a compromise between the attractive and repulsive parts of the interatomic potential. Clearly, when considering the interaction of ions with energies much higher than the thermal energies of atoms in solids, the impinging particles will achieve much shorter distances than the equilibrium positions between atoms in the solid. Therefore, a good knowledge of the interatomic potential for short distances is needed. Several ranges can be considered, depending on the distances between the nuclei. When the two atoms are far from each other (i.e., for distances r larger than the equilibrium distances between neighbouring atoms rQ) and they start to get closer, chemistry (bond formation) and van der Waals attractive forces may play a role. On the other hand, when they approach distances shorter than the Bohr radius a0 (a0 - 0.053 nm), the Coulomb potential Vcouiomb(r) between the two nuclei dominates: VCou,omb(r) = ^
-
(1.4)
r At intermediate distances (a0< r < r0) the interatomic potential is usually described by a screened Coulomb potential given by
8
Low ENERGY ION ASSISTED FILM GROWTH
V(r) =
Mi?Lx(r)
(L5)
r whereby %{r) is the screening function. For large distances (i.e., r » r0), x(r) should tend to zero, while for short distances (i.e., r « a0), %(r) should tend to unity. There are several models proposed in the literature for the screening function %{r). Among them, those proposed by Thomas-Fermi, Bohr, Moliere, Sommerfeld, Lenz-Jensen and Zeigler-Biersack-Littmark (ZBL) (Torrens, 1972; Ziegler et al., 1985; Nastasi et al., 1996) are worth mentioning. The screening function is usually expressed as a function of a reduced distance x - rla, where a is the screening length for the interaction. The value of a is an estimation of the spatial extent of the screening function. Details of the approximations used for deriving these models are out of the scope of this book and they can be found in the literature (Torrens, 1972; Eckstein, 1991). As an example, the analytical form of the screening function proposed by Moliere (which is an approximation to the Thomas-Fermi statistical model) is given by XMoiiere(x) = 0.35exp(-0.3x)+0.55exp(-1.2x) + 0.1exp(-6x)
(1.6)
In Figure 1.3 this interatomic screening function is compared with those derived from other models. 1.0 K c •2 u C
0.6
*" _g
04
|
0.2
u u
-ZBL Moliere Sommerfeld
0.8
0.0 0
1 2 3 4 5 reduced interatomic distance x
6
Figure 1.3. Different sources of screening function %(x) versus the reduced distance x = r/a (see text)
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
9
Several expressions for the screening length a have been suggested in the literature (Nastasi et al., 1996). In general, they do not differ much from each other. Within this book we will consider the expression proposed within the ThomasFermi statistical model that takes the form 0.0469
(1.7)
23
{z / +z22/3f
where a is expressed in nm. Figure 1.4 shows the screening length a in the case of O, Ar and Kr colliding ions with other atoms. Note that the screening length depends on both colliding atoms involved in the collision process. It is apparent that a decreases for increasing atomic numbers. In addition, a is roughly proportional to Z 1/3 . The interatomic screening length takes values in the range of ~10"2 nm, for which the strength of the interatomic potential between the colliding atoms is -35% of the corresponding Coulomb potential. This parameter gives an idea of the spatial extent of the screening effects in interatomic collisions.
0.020 •g" 0.018 "I" 0.016
V v
\
M 0.014 B "M 0.012 s 3> 0.010 :
Ar
•»»
/
'
-
-
-
-
•
- ^
•
Kr 0.008
•
0
I
10
'
1
20
•
—r
1
30
40
50
60
70 80
Figure 1.4. Screening lengths a for O, Ar and Kr ions interacting with other atoms with atomic number Z?, according to Eq. (1.7).
10
Low ENERGY ION ASSISTED FILM GROWTH
1.2.3. Power law approximations to the interatomic potential In order to obtain simple analytical expressions of the interatomic potential in an ion-target atom collision, the screening function %(x) can be approximated by an inverse power law of the type X(x) = X(r/a)°c rr\'-.
(1.8)
\aJ
so that the potential V(x) is proportional to x'Um. The power law parameter m takes values m = 1, 1/2, 1/3, etc., depending on the particular interaction between the colliding atoms considered. A representation of the power law potential is shown in Figure 1.5. Note how each m value in Eq. (1.8) corresponds to a fit of the screening potential in different range of reduced lengths. Thus, for example, the screening potential will be better described for reduced distances between 1 to 10, (i.e., for distances -0.01-0.1 nm) in the case of considering a power potential with m = 1/3. As we will see later, this is the appropriate range of distances for the energies involved during ion bombardment in IAD deposition of thin films. ^^^•i
-
'
"
-•
•
"
'I
%
:
'
% ^**ft» 0.1 -
% °01 1E-3
•
^ \c \r-..m=l/2 \\m=l/3
\ *x \\
:
\ 1E-4 0.1
.1
1
10
Reduced distance x Figure 1.5. Illustration of the power law approach for the screening function xM in an interatomic potential according to Eq. (1.8).
One advantage of these power law approximations to the interatomic potential is that they provide a way of analytically determining many magnitudes
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
11
related to the interaction between energetic particles and solids. However, it is necessary to identify a priori which m value in Eq. (1.8) best describes the interaction of a given type of collisions. In order to characterise energetically a collision between a moving ion with kinetic energy E, mass Mt and charge Z/ and a rest atom (M2, Z2), the reduced energy e is defined as £ =
0.69a M2 Z,Z2 M,+M2
(1.9)
where a is the screening length introduced in Eq. (1.7) expressed in nm and E is given in eV. Figure 1.6 shows the values of the reduced energy e for different combinations of ion-target atom systems as a function of the kinetic energy of the impinging ion. Note that the reduced energy e is proportional to the kinetic energy of the impinging ion, and it decreases when heavier colliding atoms are considered.
>> u
c
0.1-
0.01
•v u
s 1E-3U f
100
i
i i i i
1000
i i i i 11
10000
kinetic energy E (eV) Figure 1.6. Reduced energy e for several colliding atom systems as a function of the kinetic energy E of the impinging ion, according to Eq. (1.9).
Physically, the reduced energy e gives a measure of the energy involved in the collision and how close the ions get to the nucleus of the target atoms. In fact, the reduced energy e is inversely proportional to the distance of closest approach in the interaction of the colliding atoms dc, so that dc = ale. Note that the concept of
12
Low ENERGY ION ASSISTED FILM GROWTH
reduced energy of the collision allows an energetic classification of the atomic collisions not only by the kinetic energy of the impinging ions but also for the distances of closest approach between the colliding atoms. In practice, depending on the e value for the collision, the recommended values of the power law parameter m, corresponding to the description of the screening of the interatomic potential in Eq. (1.8), are the following (Winterbon et al., 1970): m - 1/3 for e < 0.2 (typical for Ion Assisted Deposition) m - 111 for 0.08 < £ < 2 (typical in ion implantation) m- 1 for £> 10 (typical for Rutherford scattering) According to Figure 1.6, m— 1/3 is a good approach for the range of energies typically considered in IAD. Under these conditions, the interatomic potential within the power law approach is best described for interatomic distances between -0.01 and -0.3 nm. Note that m = 1/2 may be considered as a good description of the interatomic potentials for the typical energy range of ion implantation where ion kinetic energies of several tens of keV are considered. On the other hand, m - 1 is the standard value for light ions in the MeV region as in Rutherford Backscattering Spectroscopy.
1.3. Basic concepts in classical dynamics of binary elastic collisions The interaction of energetic ions with a solid, as it takes place in IAD processes, has two main consequences from the point of view of the impinging ions, namely, a deviation from the initial trajectory and a slowing down of the ions until they stop. The first effect is mainly due to elastic collisions with the atoms of the solid (nuclear scattering), while the slowing down is due to two sources of energy transfer from the moving ion to the solid, either by the displacement of the solid atoms from their equilibrium positions through elastic scattering events (i.e., nuclear scattering again) or by the production of electronic excitations (electronic scattering). The relative importance of each type of scattering depends mainly on the energy involved in the collision. In this section, we will introduce some basic concepts needed to describe scattering processes between colliding particles.
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
13
To describe scattering processes between energetic ions and the atoms of a solid, the classical scattering theory can be applied. Within this theoretical framework, only binary two-atom collisions are considered for which classical dynamics is applied. Besides, it is assumed that electronic excitations or ionisation only enters as a source of energy loss. We will also assume that one of the two colliding atoms is initially at rest. Under these conditions, a binary collision between a moving ion and a rest atom is depicted in Figure 1.7. Considering energy and momentum conservation laws in a binary collision, the ratio of the projectile energies before, E0, and after, Eh the collision event is given by
^L-
M, cos9 ± (M22 - M] sin2 o)' M,+M2
(1-10)
where 6 is the scattering angle. The positive sign in Eq. (1.10) holds if M ; < M2. In this case any scattering angle 6 is possible. However, if M; > M2, the maximum angle through which M; can be scattered is arcsin{M2IMi).
before collision event ^\Eo,v0 M7)-2-V
f \ At rest (M.
E„v,
after
collision
E2>*2 Figure 1.7. Scheme of a binary collision between a moving ion (Mi) and a rest atom (M2).
The energy transferred in a single collision to the rest atom, E2, can be calculated by the expression
14
Low ENERGY ION ASSISTED FILM GROWTH
E2
-A- = -/
E0
4M,M2
2
^ r C O S
(1.11)
(M.+MJ
where the dispersion angle 0 is defined in Figure 1.7. Thus, the maximum energy TM transferred in a head-on collision is 4M M
Eq. (1.12) establishes a first limitation regarding the ability to transfer energy in single collisions processes. Thus, the maximum energy transferred in a single collision process (i.e., TM/E0 = 1) can be only achieved if Mj = M2, while if the two masses are different, TM/E0 will be always less than unity. It is also important to note that the energy transferred in a single collision does not depend on the charge of the colliding particles. According to Eq. (1.12), there will be a minimum kinetic energy E,h for a primary ion to produce an atom displacement of energy Ed. This energy is given by
*
4M,M2
"
Thus, if the ions have a kinetic energy E0 < Eth they can be only stopped within the target lattice by incorporation in interstitial sites, i.e., they will not create Frenkel pairs.
1.3.1. Ion energy loss rate Up to this point, only some basic ideas have been introduced to describe the elastic collision events. Now, it is necessary to apply the previous concepts to determine parameters that are directly related to the processes taking place in practical IAD film growth. Thus, we are concerned with the distribution of the "implanted" ions and how their energy is shared with the other atoms of a solid surface. When an energetic ion penetrates a solid, it undergoes a series of collisions with the atoms and electrons of the target. These collisions result in ion energy
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
15
losses along the ion path. Two main consequences arise from the slowing down of the primary ions: the transfer of energy and momentum to the atoms of the target and the implantation of the impinging ions. As a good approximation, the energy loss rate of an ion due to the interaction with the solid target dE/dr, also called the stopping power, can be expressed as
dr
dr
dE + — dr
(1.14)
where the subscripts n and e denote nuclear and electronic collisions, respectively. Nuclear energy losses show up as atom displacement and electronic energy losses as electronic excitations and ionisation processes. In the previous sections, only nuclear stopping was considered. However, we will see that, although small for the range of energies considered in IAD processes, there is a percentage of the primary kinetic energy of the impinging ions that is used in producing electronic excitations, and not atom displacement. Nuclear collisions can involve large energy losses and significant angular deflection of the trajectory of the ions. These processes are also responsible for the production of lattice disorder. Electronic collisions involve much smaller energy losses per collision, negligible deflection of the ion trajectory, and negligible lattice disorder. The relative importance of these two energy loss mechanisms depends on the kinetic energy E0 of the impinging ions, and the particular characteristics of the colliding ions (mainly through Z/ and Z2). We can advance that nuclear stopping predominates for low E0 (few hundreds of eV) and high Zh whereas electronic stopping takes over for high E0 (several hundreds of keV) and low Zh Typical units for the energy loss rate are eV/nm. Usually the stopping power is expressed as a function of the corresponding nuclear and electronic stopping cross sections, S„(E) and Se{E) as
dr
•NSnJE)
(1.15)
where N is the atomic density of the target and E the kinetic energy of the energetic ion. The stopping cross sections are usually expressed as a function of the reduced energy introduced in Eq. (1.9). They are related by the expression
16
Low ENERGY ION ASSISTED FILM GROWTH
SnJE)
M = 18.1aZlZ2-—±—SJe) M, +M2
(1.16)
This reduced notation enables an easy comparison between collisions with different energy regimes and also to compare electronic and nuclear losses. In the following we will introduce expressions to evaluate stopping powers and stopping cross sections. They will allow us to estimate analytically different magnitudes of interest in IAD processes.
1.3.1.1. Nuclear stopping To characterise in energy terms a scattering event it is necessary to know not only the particular energy transferred in a single collision event, but also the distribution of energy losses, i.e., the probability for a particular energy loss E2 to take place (in the following we use T = E2 for clarity). This probability is given by the energy transfer differential cross section da„(E)/dT. Without going into details, and making use of the power law approximation (i.e., for interatomic potentials of the form V(r) ~ fVm as introduced in section 1.2.3.), the energy transfer differential cross section can be expressed as da„(E)= dT
C, EmT
a-17)
where Cm is a constant that depends on the type of colliding atoms. It is given by
C M =~A m a
2
2Z,Z2e2
\2m
(*,M, \ M22 , V
(1.18)
J
where ^ is also a constant related to the type of interatomic potential used within the power law approximation. In the case of the Thomas-Fermi statistical potential (cf. section 1.2.2), X„ takes the values given in Table 1.1. It was mentioned in section 1.2.3. that the energy range of interest in IAD is better described with m = 1/3, for which Xm = 1.309. Note also that for m - 1, we recover the well known Rutherford scattering cross section applied for high energy
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
scattering processes such as those taking place in standard Backscattering Spectroscopy (RBS).
17
Rutherford
Table 1.1. Values taken by Am within the power law approximation to the Thomas-Fermi statistical model (Sigmund, 1981).
m
4m
1 0.5
0.5 0.327
0.333 1.309
0.191 2.92
0.055
15
0 24
By nuclear stopping, we refer to the average energy loss that results from elastic collisions between the impinging ions and the target atoms. Nuclear collisions can involve large energy losses and significant angular deflection of the ion trajectory. This process is responsible for the production of lattice disorder. The nuclear stopping cross section is related to the energy transfer differential cross section introduced before (cf. Eq. (1.17)) as Sn(E) = £»TCi°^dT
(1.19)
where TM and Tmin are the maximum and minimum energy transferred in the scattering processes. TM was introduced in Eq. (1.12) and Tmin is usually referred as the energy needed to displace an atom from its lattice site, i.e., the displacement energy Ed introduced in section 1.2.1. Using the power law representation and Tmin = 0, the nuclear stopping cross section as a function of the reduced energy Sn(e) can be calculated by the expression
2( 1 - m) The previous equations (1.17-1.20) allow us to calculate energy loss rates and nuclear stopping cross sections based on the Thomas-Fermi model with an accuracy of -20% (Nastasi et al., 1996). For the range of energies used in IAD (i.e., for £ < 0.2, m = 1/3), the nuclear stopping cross section and the corresponding rate for energy loss raises as E1B while for higher impinging energies as those used in implantation processes (i.e., for 0.1 < e < 2, m = 1/2), Sn is independent of E. In the case of considering high-energy interaction, as in RBS characterisation, for example, (e > 10) we recover the Ex dependence typical of Rutherford scattering.
18
Low ENERGY ION ASSISTED FILM GROWTH
Figure 1.8 shows the nuclear scattering cross sections for the different energetic regions considered in this section according to power law potentials (i.e., different m values) compared to that obtained from a full calculation within the Thomas-Fermi model. The electronic stopping cross section, also reported in this figure, will be discussed in next section. 0.5
Se(Jfc = 0.4) S (k = 0.2)
'to 1
0.4-
CO *
0.3
"to 00
1
0.2S (Thomas-Fermi)
0.1
0 . 0 ~ **"" i
1E-3
i i i 11111
0.01
i
i
i 111 i i |
0.1
i
i
i 111111
1
i
i i 111n|
10
reduced energy e Figure 1.8. Reduced nuclear and electronic stopping cross sections versus the reduced energy e, according to Eqs. (1.20) and (1.21).
Note for example that the approach of considering power law interatomic potentials with m = 1/3 is fully justified because in most collisions taking place in IAD processes, £ is smaller than 0.1 (compare the results shown in Figures 1.6 and 1.8). As an example, to get an idea of the importance of the nuclear stopping in IAD thin films, the stopping power calculated according to Eqs. (1.9), (1.16) and (1.20) for 400 eV Ar ions impinging on a silicon substrate is -270 eV/nm (i.e., the impinging ions lose 67% of their original energy in a path of 1 nm). This value indicates that a very shallow penetration of the ions is expected for the energy ranges typical of IAD processes. However, the description of the ion energy losses due to elastic collisions with the target atoms considered in this chapter within a power law approximation, is a very simplified image of the process. There are, in the literature, sophisticated models based on more complete and realistic descriptions of the particular
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
19
interaction between colliding ions (Ziegler 1985). Our purpose here is only to give simple analytical expressions that enable a fast evaluation of the tendencies expected when changing experimental parameters such as the atomic number of the colliding atoms or the kinetic energy of the impinging ions. Readers interested in a more accurate description of the nuclear stopping cross section are referred to those publications.
1.3.1.2. Electronic stopping In addition to nuclear stopping, it is also important to have estimations of the electronic stopping power, i.e., the energy spent in electronic excitations that do not contribute to atom relocation. Among the different models available in the literature that take into account electronic losses (Nastasi et al., 1996), we will consider here the extensively used Lindhard-Scharff model (Sugiyama 1981), in which the reduced electronic stopping cross section in the low energy regime of interest in IAD can be expressed as Se(e) = ke'/2
(1.21)
where k is the so-called electronic energy-loss parameter given by
72/37l/2
Z,;
k=
Z,2
(
7+ ^ 1
12.6M'2/2(Z2,'3 + Z2/3
f*
(1.22)
In practice, in most cases k takes values between 0.2 and 0.5 (see Figure 1.9). Note that the contribution of the electronic stopping is more important when light energetic ions interact with heavy target atoms. In general, in the low energy regime (i.e., for £ < 0.2) the energy involved in electronic excitations is small compared with the energy spent in nuclear elastic collisions. This is shown in Figure 1.8 where the nuclear and electronic contributions to the stopping cross section are compared. However, since the electronic loss rate increases with the energy of the ion, for collisions characterised by reduced energies £ of the order of 1, both nuclear and electron cross sections have similar values. In these conditions, the electronic stopping may easily account
20
Low ENERGY ION ASSISTED FILM GROWTH
for a considerable amount of the energy otherwise available for atom displacement during IAD processes.
0.8
o• >> • i * 61) U >at s S a © u
*-•
Dfi
0.4
08
t-
a a
0.2
a*
01
o.o-i—.—,—.—i—.—i—.—i—.—,—.—i—.—i—.—I 10 20 30 40 50 60 70 80
Figure 1.9. Electronic energy-loss parameter k for O, Ar and Kr ions, as a function of the atomic number Z2 of the target atoms, according to Eq. (1.22). M2 = 2.2xZ2 has been considered for simplicity.
Following the example of 400 eV Ar bombardment of a Si target mentioned in the previous section, the expected electronic energy loss rate according to Eqs. (1.16), (1.21) and (1.22) is -11 eV/nm, i.e., much less than the -270 eV/nm employed in nuclear elastic collisions.
1.4. Range of energetic ions in solids Until now, we have introduced expressions to estimate the rate at which the kinetic energy of an impinging ion is lost either in electronic excitations or through elastic collisions with the target atoms. As a result of this energy loss, the impinging ion will slow down until it stops completely. The total path R travelled in a solid by the penetrating ion before it stops is known as the ion range. The range that an ion travels in a solid can be calculated from the energy loss rate as
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
R=C°
(dE)
dE
21
(123)
J
E<> y dr
j
The main parameters governing the range R are the kinetic energy E of the ion, and the atomic numbers of the ion and atoms in the target. In fact, if we neglect the electronic stopping and make use of the power law approximation, R can be evaluated from Eqs. (1.15), (1,16), (1.20), and (1.23) by the expression R =
0.0796(M,+M2)2 Na2 M,M2
l-mc2m rnlm
where N is the atom density of the target in atoms per nm3, a the interatomic screening distance defined in section 1.2.2, and m and Am are parameters in the power law approximation of the interatomic potential (cf. Table 1.1). Figure 1.10 shows the range R that can be obtained for Ar and O ions travelling in Si and Zr solids as a function of their kinetic energy, according to Eq. (1.24). Note that for a fixed substrate, the heavier the impinging ion, the higher the range R, while for a fixed impinging ion, an increase on the atomic number of the atoms of the substrate Z2 results in an increase of the ion range R. This behaviour is related to the ability of transferring energy in a collision event, and it follows roughly the same dependence as the maximum energy transferred in a single collision event TM introduced in Eq. (1.12). Another conclusion from Eq. (1.24) is that for standard IAD processes (i.e., m = 1/3), R varies proportional to EVi'. Note also that electronic stopping will tend to decrease the actual ranges calculated with Eq. (1.24). It is also worth noting that the range R for low energy ion beams (E < 100 eV) takes values less or around that of 1-2 nm, i.e., a very shallow penetration is expected in these cases, typical of IAD processes. The time t for a primary ion with initial energy E to come to rest can be approached by the ratio of its range R divided by half of its initial velocity, (i.e., t ~ 2Rlv). This gives a time scale of ~10"13 s for the process, i.e., of the order of the lattice vibrations. Then, the approach made previously (cf. section 1.2.2.) that the atoms of the solid are at rest during the interaction with the energetic ions is fully justified.
22
Low ENERGY ION ASSISTED FILM GROWTH
S c M
a U
i i i 11
1000
10000
kinetic energy E (eV) Figure 1.10. Typical ion ranges involving O and Ar ions impinging in Si and Zr substrates, calculated according to Eq. (1.24). Solid lines: m = 1/3; dotted lines m = 1/2.
The range R introduced above refers to the total path travelled by the ions before they stop. From a practical point of view, it is also interesting to know the socalled projected range Rp, i.e., the total path travelled by the ions along the direction of incidence. A schematic representation of the total range R and projected range Rp of penetrating ions in solids is depicted in Figure 1.11.
Figure 1.11. Schematic representation of the total ion range R and the ion projected range Rp .
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
23
The relationship between R and Rp depends again on the kinetic energy of the impinging ion E, and on the particular atomic numbers of the projectile ion and target atoms. Thus, the ratio Rp/R decreases as the mass ratio A/2/M/ increases. This effect is easily understood if we consider the expected dispersion of the ion trajectories. Heavy ions impinging in light substrates (Mi > M2) will have weaker deflection of their trajectories than light ions impinging on heavy substrates (M/ < Mi). In practice, it has been proposed the following phenomenological expression relating R and Rp (Lindhard et al., 1963)
R
1 + B(M2/M,)
where B takes values of -0.6 for m - 1/3 and -0.33 for m - 1/2. Figure 1.12 shows the ratio R/R according to Eq. (1.25). The behaviour shown in Figure 1.12 is similar to the one reported by Winterbon et al. (1970).
0.0-1 0.1
. 0.2
.— 0.5
1
. 2
.— 5
10
M2/M1 Figure 1.12. Projected range Rp and the total range R as a function of the mass ratio MJM, for m = 1/2 and m = 1/3, according to Eq. (1.25).
In the literature there are other strategies to estimate projected ranges of ions. The most popular is the Montecarlo calculation made by TRIM software (Ziegler 1985). If the reader is interested in the detailed description of TRIM code, or other sources of computer simulation of ion-solid interaction, he should consider the review by W. Eckstein (1991). To make a critical view of Rp values obtained
24
Low ENERGY ION ASSISTED FILM GROWTH
from Eqs. (1.24) and Eq. (1.25), Figure 1.13 shows a comparison between Rp obtained with the TRIM code and that with the previous Eqs. Note that the degree of agreement is in general very good although the electronic losses are not included in Eq. (1.23) and (1.24). According to Figure 1.13, the projected ranges of 0 2 + and Ar+ ions with kinetic energies of 1000 eV in a Si matrix are 2.8 and 3.1 nm, respectively. I I • • |
1
1
1..,.,
,,,l
—,
s •.
I
I
I i-TTiri
r-1 i
i i i ii
10-
10-
s/
ii
,—,—,,1—T
j/gjyX?
4
'•
/jJ^Xfr
U-
o— s\A
p^Ar— Si
1-
100
1000
E (eV)
10000
100
1000
10000
E (eV)
Figure 1.13. Comparison of the projected ranges calculated according to the power law approximations (solid lines) and TRIM simulations (symbols) for several ion —» target systems.
Until now, we have always dealt with monoatomic target materials. In the case of IAD growth of compounds as oxides, nitrides or carbides this is not the case. For a compound AxBy (an oxide for example), it has been proposed a projected range given by
Rp(AxBy)=N
(Rp(A)/NA)(Rp(B)/NB) (yRJA)/NA) + (xRJB)/NB)
(1.26)
where x + y - 1, Rp(A), Rp(B), NA, and NB are the projected ranges and the atomic densities in pure A and B targets respectively, and AT the atomic density of the compound. Another aspect refers to the uncertainty in the determination of Rp. The trajectory of energetic ions in condensed matter is a stochastic process. The
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
25
projected range Rp introduced before is the most probable value among a distribution of distances for an ion to come to rest. The standard deviation from the mean distance Rp is called the range straggling ARP. This straggling is expected to be small for heavy ions impinging in light targets (i.e., M2I' Mj< 1), but it can increase dramatically as the ratio M2 I M; increases, i.e., when light ions bombard heavy substrates (Winterbon et al., 1970). In many cases, it is a reasonable approach to consider that ARP = 0ARp (Nastasi et al., 1996).
1.5. Spatial distribution of deposited energy As a result of the slowing down of the impinging ions, their initial kinetic energy is deposited along the ion track within the film. The depth profile of the deposited energy is not necessarily identical to the implantation profile. The energy loss rate depending on the position of the ion in its track x is known as the deposited energy depth distribution function FD(x). If we neglect electronic stopping and the atomic displacement threshold, FD(x) can be obtained within the power law approximation by (Winterbon et al., 1970)
FD(x) = ^-{l-X-Jm" D
2mR{
(1.27)
R)
Since Eq. (1.27) does not take into account the displacement threshold, it really refers to the spatial distribution of the energy available to displace atoms. In the case of an IAD process with high-energy ion bombardment (i.e., m - 1/2), FD(x) takes a constant value EJR, i.e., the deposited energy is evenly distributed along the range R. On the other hand, in the more common IAD case of medium and low energy bombardment (i.e., m = 1/3), FD(x) is proportional to (1 - x I R)05, i.e., the energy of the impinging ion is lost preferentially at the beginning of the ion track. This means that in typical IAD processes, the deposited energy is located in the shallower regions as compared with the penetration range of the ions, i.e., at the surface of the bombarded substrates or growing films. Figure 1.14 illustrates the variation of FD(x) with the primary energy of the impinging ions for the cases m 1/2 and m = 1/3. Strictly speaking, previous Eq. (1.27) is only valid for ions slowing down continuously along a straight line. This approach breaks down for M, < M2. Taking
26
Low ENERGY ION ASSISTED FILM GROWTH
1.SE/R^ \ ^ m = l/3 m = l/2 fc."
0.5E/R-
—i
0.0
o.a
0.4
0.6
1
1
0.8
—'
1.0
x/R Figure 1.14. Deposited energy depth distribution function FD(X), i.e., location of the deposited energy versus path travelled by the ion, for high (m = 1/2) and low energy IAD processes (m = 1/3).
into account more realistic trajectories for the ions, Winterbon et al. (1970) found that the mean location of the deposited energy or average depth of damage {FD(x)) = (X) depends on the M21 Mt ratio as it is shown in Figure 1.15. Note that the implantation profile Rp versus R does not coincide with the damage profile (X) versus R, (cf., Figures 1.12 and 1.15). The straggling of the energy profile in the direction parallel (AX 2)m and perpendicular (Y2)m to the ion trajectory are also depicted in Figure 1.15 for the typical energies involved in IAD processes (m = 1/3)
Ma/M, Figure 1.15. Average depth of damage (X) and damage straggling in the direction parallel
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
27
and the higher energy regime (m = 1/2).
1.6. Damage induced by ion bombardment As a consequence of the energy deposition by ion bombardment, damage will be produced in the solid target typically as atom displacements. Therefore, only nuclear elastic scattering is responsible for damage effects in solids. In the following, the mechanisms through which the energy of the impinging particles is deposited in the target will be described to quantify this damage production.
1.6.1. Primary knock-on atoms formation Before an energetic ion comes to rest in a solid, it makes a number of collisions with the lattice atoms. In these collisions, some energy can be transferred from the ion to the target atoms which, as a consequence, are displaced from their lattice sites. Lattice atoms that are displaced by incident ions are called primary knock-on atoms (PKA). The PKAs can in turn displace other atoms usually considered secondary, tertiary knock-on atoms and so on, thus creating a cascade of collisions or displacement cascades. For simplicity, every displaced atom is considered a PICA. Figure 1.16 illustrates schematically the formation of these displacement cascades of PKAs. 9
incident particle
© primary knock-on atom (PKA) 0 secondary knock-on atom (PKA)
. A
$ tertiary knock-on atom (PKA) gl vacancy
Figure 1.16. Scheme of a cascade formation of PKAs
28
Low ENERGY ION ASSISTED FILM GROWTH
The energy required to displace an atom from its lattice site, the so-called displacement energy, was introduced previously in section 1.2.1. It should be noted that if the energy transferred is smaller than the minimum displacement energy (e.g. -20 eV), there will be no atom relocations and therefore no damage. Primary ions and PKAs will lose energy through both electronic and nuclear collisions, but only the latter causes lattice disorder around the ion track and is responsible for the production of structural damage. This means that not all the primary kinetic energy of an impinging ion is converted into damage. The energy that is not lost in electronic excitations is known as the damage energy v(E). It can be estimated analytically within the Norgett-Robinson-Torrens (NRT) model of damage energy (Norgett et al., 1975). Thus, according to this model, the damage efficiency defined as the ratio between the damage energy and the total energy of the ion v(E)/E is given by
= [l + *(£ + 0.0424£075 +3.40l£ 01667 )]~'
^
(1.28)
E where k is the electronic energy-loss parameter (cf. Eq. (1.22)). Figure 1.17 shows the dependence of the damage efficiency v(E)/E as a function of the initial energy of the particle E for different ion-target systems According to Figure 1.17, the damage efficiency decreases as the ion kinetic energy increases. Additionally, note that for typical IAD conditions v(E) accounts for the -90% of the initial energy of the ion. However, when heavy ions bombarding light substrates are considered, the energy involved in electronic excitations increases considerably as is also shown in Figure 1.17. Thus, the damage efficiency of Ar bombardment on a Si matrix is -0.8 for 100 eV kinetic energy of the impinging ions, decreasing to -0.7 at 2000 eV. Eq. (1.28) only applies for pure elemental solids but may be used to have a rough estimation of damage efficiency in compounds. A detailed estimation of damage energy in materials with more than one type of atom in their structure require more complicated models, which are beyond the scope of this book (Nastasi et al., 1996).
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
I
I I I I I |""
I
I
|'"-c I I I I
1000
100
29
10000
E (eV) Figure 1.17. Damage efficiency x>(E)/E as a function of E for different ion-target systems according to Eq.(l .28)
Another interesting magnitude to quantify the damage induced by particle bombardment is the average number of displaced atoms in a cascade produced by a primary ion (or by a PKA) of energy E. Usually this magnitude is known as the displacement damage function (Nd (E)). For ion kinetic energies lower than the displacement energy, there is no possibility for atom displacement (i.e., (Nd (E)) = 0, if E < Ed). At the same time, as long as the kinetic energy of the ion is greater than Ed, and smaller than 2E/1;, (where 2; accounts for energy losses due to electronic excitations; in practice t, - 0.8-0.9) then (Nd (E)) = 1. On the other hand, for higher energies {Nd (E)) = %v(E)/2Ed, so that in general
(Nd(E))=
0 1 $v(E)/2Ed
(for
0<E<Ed)
(for
Ed<E<2Ejt;)
(for
2Ed/$<E«~)
(1.29)
{Nd (E)) provides an easy evaluation of the number of displaced atoms for a particular IAD condition. Thus, for example, considering a typical IAD process with Ar+ ions of 400 eV and a particular substrate with an average displacement energy of -40 eV, there will be ~4 displaced atoms due to the interaction of each impinging Ar ion with the target atoms.
30
Low ENERGY ION ASSISTED FILM GROWTH
1.6.2. Spike formation As we have seen in previous sections, when a low energy ion penetrates a solid, it slows down mainly due to nuclear elastic collisions with the target atoms. As a consequence, the ion will produce PKAs along its path (i.e., along the range R). The mean free path between successful collisions in the production of PKAs XD will decrease until a PICA is generated in every lattice site along the ion path. In these conditions a large number of point defect (atom displacements) will be created in a short time, as is schematically illustrated in Figure 1.18. The concept of spike in irradiated materials is defined as a high-density cascade of collisions that have been created in a limited volume where the majority of the atoms are temporarily in motion.
ooo o o\roooooooooo o oooo \ o*o o% o o o o o oooo \ oooooo o o ooooo \ o o% o o o o o o o»o o o o ^ c^o o o o o o o o o o*o o o o CLO o o o o o o o o o o o o o aoo o o o o o o ooooooooooooooooo ooooooooooooooooo ooooooooooooooooo ooooooooooooooooo Figure 1.18. Scheme of a spike formation at the surface of a solid target by an energetic ion.
According to the power law approximation, the average path length XD(E,T0) of a given ion with kinetic energy E with a transfer of energy per collision greater than T0, can be obtained from the equation Xd(E,T0) = ~mL
1M1
°
NCJTX-T?)
(1.30)
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
31
where m is the parameter describing the screening of the interatomic potential (cf. section 1.2.3.), Cm was introduced in Eq. (1.18), TM given by Eq. (1.12) is the maximum transferred energy, and N the atomic density in atoms per nm . When Xd is of the order of the interatomic distance, a PKA is generated at every lattice site. Figure 1.19 shows the values obtained from Eq. (1.30) in the case of low energy bombardment (m = 1/3) for Ar ions impinging an Si target. According to this figure, for typical IAD conditions we will find spike behaviour almost throughout the whole range of the impinging ions. This means that the spike regime will be localised just at the outermost surface layers of the target.
•
1
1
1
50
r——r
1
1 — |
100
1
200
300
Figure 1.19. Mean free path Xj between collisions with an energy transferred greater than T0 between recoils of Ar ions of various energies £ in a Si target, according to Eq. (1.30) in the case of low energy bombardment (m = 1/3).
It is also apparent in this figure that as the ion energy decreases, the distance between recoils becomes smaller, until Xd approaches the interatomic spacing between atoms in the target. Then a PKA will be produced in every lattice site along the ion path. As a consequence, a highly damaged region is formed, producing a volume of material that is composed of a core of vacancies surrounded by a shell of interstitial atoms, as is schematically depicted in Figure 1.18. This highly damaged volume of material is known as displacement spike.
32
Low ENERGY ION ASSISTED FILM GROWTH
1.6.3. Thermal spikes When the formation of a spike comes to an end, all the moving displaced atoms reach a point at which they have insufficient energy to cause further displacements. The neighbouring atoms will share the energy as vibrations or heat. This period of lattice heating is known as the thermal spike phase of the collision cascade and it may exist for several picoseconds before being quenched to ambient temperature at a rate of 1010-1012 K/s. The quench time for a thermal spike should increase with the size of the spike and it is inversely proportional to the thermal diffusivity of the target. The time scale for this process will again be of the order of magnitude of lattice vibrations , i.e., 10"12-10"13 s . The thermalisation process taking place after the deposited ion energy in the spike can be described considering the solid as a continuum. The evolution of the temperature T(r,t) in the proximity of the spike region with the distance r and the time t can be considered as following the law (Hofsass et al., 1998) T(r,t)oc—^—
(Dtf
-r' ' 4Dt
(1.31)
where Q is the thermal energy deposited in the spike, D is the diffusion coefficient and n is a constant (1 < n < 1.5) that depends on the shape of the cascade. Eq. (1.31) is suitable to describe qualitatively the behaviour of a thermal spike, though it may fail for diffusion lengths of the order of, or smaller than the interatomic distances. Typical values for D are ~10 n nm2/s and the characteristic relaxation times t ~10"12 seconds. Figure 1.20 shows the evolution of the temperature during the thermal spike phase as a function of the distance to the spike location during the next picoseconds after the energy absorption, according to Eq. (1.31). Note how there will be a thermal wave due to the thermal diffusion that will raise the local temperature for few picoseconds. This thermal wave may also be responsible for damage production.
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
33
s
i
I Distance (nm)
Figure 1.20. Evolution of the temperature during the thermal spike phase as a function of the distance to the spike location. Each curve corresponds to the temperature profile 0.5, 1, 2 and 4 ps after the energy absorption, according to Eq. (1.31) .
1.6.4. Density of the deposited energy Another interesting magnitude is the mean damage energy (6D) per atom in a cascade. (6D) gives an idea of the local temperature reached in the spike regions, before thermal relaxation. A rough estimation of (6D) can be obtained from the expression (Nastasi et al., 1996)
{eD)=^^ x
'
(i.32) NV cas
where v(E) is the damage energy defined in Eq. (1.28), N is the atomic density of the target and Vcas the volume of the central part of a cascade. The 0.32 factor accounts for the statistical fraction of damage energy residing within one standard deviation of the mean volume of the cascade where the energy is deposited. As a first approach, (6D) is expected to be proportional to E'~6m (note that v(E) is proportional to the kinetic energy of the impinging particles E, and Vcas should be
34
Low ENERGY ION ASSISTED FILM GROWTH
proportional to R3 and therefore to E6"1). Thus, under IAD conditions (i.e., m = 1/3), (6D) should be proportional to Ex. To evaluate the volume of the central core of an ensemble of cascades Vcas, there are different approaches to be considered. If we assume a spheroid shape for the ensemble of cascades it has been proposed (Nastasi et al., 1996) V^S'^Axfy)
(1.33)
where 8 is a correction factor and the rest is the volume of a spheroid defined by the straggling of the energy deposited in the direction parallel (AX2)1/2 and perpendicular (y2)m to the ion trajectory (cf. section 1.5). The correction factor S accounts for the different types of cascade formation, represented in Figure 1.21. Thus, the energy deposited when a light ion impinges a heavy substrate will be localised close to the ion track (represented in the figure by full lines). However, when heavy projectiles interact with light elements, highly energetic PKAs can be formed, thus producing sub cascades that will distribute the energy far from the primary ion track. 8 depends on the mass ratio M2/Mj in a manner very similar to Rp/R, so that we can consider <5 ~ Rp/R. For example, in the case of Ar+ ions with 400 eV kinetic energy impinging on an Si surface, the volume of the central part of the cascade Vcas can be approached by -0.3/?3 where R is the corresponding ion range given by Eq. (1.24). R takes the value of -2.1 nm in our case, so that Vcas~0.3 nm3. Thus, the energy deposited per Si atom in the cascade volume (0D) is ~7 eV. There are, in the literature, other models to describe the volume of a spike. One possibility is to consider a cylindrical volume, whose length is related to the range of the energetic particle and its section mainly dependent on the energy of the particle (Hofsass et al., 1998). Considering that the deposited energy follows a Maxwell-Boltzmann distribution, the mean deposited damage energy is related to the local temperature Tlocal a s
K)
= \Wiocai
= 1.29xlO-4Tlocal
(1.34)
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
Heavy ion impinging in a light substrate
35
Light ion impinging in a heavy substrate
Figure 1.21. Different shapes for cascades formation. Thick and thin lines represent the ion track of the impinging ion and the track of PKAs created. The dots represent displaced atoms.
Typical (0D) are in the range of few eV/atom, as it is shown in the previous example of 400 eV Ar+ ion impinging on Si. This corresponds to local temperatures at the cascade site of the order ~104 K far above the melting point of the corresponding solids.
1.7. Sputtering As a consequence of ion bombardment, a series of phenomena takes place on the target. Some of the atoms located at the outer layers of the bombarded target may be ejected out of it, i.e., sputtered by the interaction with the impinging ions. Another effect is the ion mixing of different atom species present in the target or the ion mixing in buried interfaces when they are reached by the impinging ions. Besides, the localised high temperatures achieved at every ion impact site may induce phase transformations. In the following we will describe the basic concepts related to the sputtering phenomena, while ion mixing and phase transformations will be dealt with in Chapter 3.
36
Low ENERGY ION ASSISTED FILM GROWTH
We can describe the sputtering process phenomenologically as follows. Bombarding ions transfer energy in collisions to target atoms, which recoil with sufficient energy to generate other recoils. Some of these backwards recoils will approach the surface with enough energy to escape from the solid. These secondary recoils are responsible for most of the atom ejection, i.e., sputtered material. Following the power law approximation to describe interatomic potentials, the region Axsp from which the particles are ejected from the target can be approximated by (Sigmund, 1981) 1-m
E
m
Axm= sp
(1.35) l-2mNCm
where E is the kinetic energy of the impinging particles, N the atom density and Cm the material dependent parameter defined in Eq. (1.18). In the case of low energy bombardment (i.e., m = 1/3), Axsp is proportional to £° 66 . Thus if the primary energy of the impinging ions is doubled, the region from which target atoms are sputtered increases -60%. In practice, Axsp takes values between 1 to 4 monolayers, depending on the particular system under investigation.
1.7.1. Sputtering yield An important magnitude to quantify the efficiency of the sputtering phenomena is the so called sputtering yield Y, which is defined as the ratio of the mean number of emitted atoms divided by the number of incident particles. It can be estimated by (Sigmund 1981)
Y{E^j_^mm
(1.36,
E
b
where Sn(E) is the nuclear stopping cross section given by Eqs. (1.15) and (1.19), Eb is the surface binding energy of the atoms in the solid expressed in eV and afli) is a correction parameter that depends on the angle of incidence /J of the impinging ions with respect to the surface normal. In the literature there are other semi-empirical equations to evaluate the sputtering yield Y that have shown a high degree of
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
37
consistency with experimental measurements. Readers interested in more specific works are referred to the works of Ziegler et al. (1985) or Yamamura (1989). The dependence of the sputtering yield with the type of incident particles, as well as with their kinetic energy is mainly determined by the nuclear stopping power Sn(E), so that, the sputtering yield increases with the mass of the incident particles if the energy and target atoms are fixed. Figure 1.22 shows experimental results and calculated values for the sputtering yield of a Ni substrate for different incident particles. Note that the sputtering yield increases until -1000 eV and then saturates for higher kinetic energies.
Figure 1.22. Sputtering yield of several bombarding ions in Ni. Reproduced from Ziegler et al. (1985) with permission.
The surface binding energy Ebs is a material dependent parameter (see section 1.2.1.) that in practice takes values between 3 and 8 eV (Kittel, 1976). The correction parameter afli) depends on the target atom to incident ion mass ratio (i.e., M2/Mi) and the angle of incidence fi of the impinging ions with respect to the surface normal. Thus, at normal incidence a(0) ~ 0.2 for M2/Mi < 1, and then it increases continuously to a(0) ~ 0.4 and 0.6 for M2/M, equal to 3 and 7, respectively (Sigmund 1981). For angles of incidence /? <70°, afli) increases with fl due to the increase of the deposited energy near the surface. It is found that afli)/a(0) ~ (cos/J,K
38
Low ENERGY ION ASSISTED FILM GROWTH
where the exponent / again depends on the mass ratio M^/M, and takes values between 1 and 2. For angles of incidence of -75°, a(P) and hence the sputtering yield reaches a maximum. If p is increased above -75°, the sputtering yield decreases. This is due to the fact that the probability of the particle beam being reflected increases for glancing incidence. Under such conditions the average path length travelled inside the target, and hence the energy loss, decreases. To illustrate the dependence of the sputtering yield with the angle of incidence of the impinging ions, Figure 1.23 shows the result of the Montecarlo TRIM sputtering yield calculations of Ar+ and 0 2 + ions impinging on a Si0 2 target.
-i
0
•
i
20
•
1
•
40
1
60
1
r-
80
Angle of Incidence p Figure 1.23. TRIM simulations of the angular dependence of the sputtering yield of Ar and 02* ions impinging in Si02
In IAD processes, sputtering effects play a significant role with respect to the film growth. Thus, a first evidence is that the removal of material during thin film growth will decrease the growth rate. Other effects may be related to the socalled preferential sputtering effects, i.e., the preferential removal of a certain type of atoms in multielemental samples. This may have as a consequence a change in the final stoichiometry of the thin film. In this respect it is very common that the surface of ion bombarded oxide targets becomes depleted in oxygen with respect to the stoichiometric oxide
1.7.2. Angular distribution of the sputtered atoms According to the phenomenological description of the sputtering process mentioned above, the angular distribution of the ejected particles dY/dQ is rather independent
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
39
of the angle of incidence of the bombarding ions, at least for angles of incidence lower than 60°. This is because the recoil cascade loses the information of the primary projectile-target interaction during the stopping of the particles. As an example, Figure 1.24 shows the angular distribution dY/dQ of the sputtered Ti atoms as a function of the angle of emission 0e for 1000 eV Ar4" ions bombarding a Ti substrate. Two situations are compared: normal incidence of the bombarding ions (left) and bombardment with an angle of incidence 60° off normal as it is described in the inset of the figures. Note that the maximum of emitted particles is always at a direction normal to the surface and that only minor differences are observed for negative and positive values of 9e. In practice it is found that dY/dQ «= (cmBe)~n with n between 1 and 2.
llMM)eVAr + —TI \T*
60°
S i
0
15
30
45
60
75
90
(
.
.
)
i
, ^ -
r
-
T n
- - ^ -
-90 -75 -60 - 4 5 - 3 0 - 1 5
angle of emission 0
T
0
- ^ -
r 7
^ -
r n
- ^ .
r r
^ ^
r
. . , . -
r
. .
T
-
r
-
r i
15 30 45 60 75 90
angle of emission 0
Figure 1.24. Angular distribution of sputtered particles for 1000 eV Ar4* sputtering on a Ti substrate. Normal incidence (left) and angle of incidence of 60° (right).
1.7.3. Energy distribution of the sputtered atoms It is also interesting to know the energy distribution of the ejected material in the sputtering process dY/dEe. A detail description can be found in Hofer (1991). The atoms that- escape from the surface are those reaching the surface with a kinetic energy higher than the surface binding energy and with momentum pointing away
40
Low ENERGY ION ASSISTED FILM GROWTH
from the material. Montecarlo simulations using the TRIM code allow us to evaluate it. Figure 1.25 shows the energy distribution of Ti atoms corresponding to 1000 eV Ar+ bombardment with normal incidence into a Ti substrate. It has been found experimentally that dY/dEe decreases as the inverse of the square of the kinetic energy of the ejected atoms.
1000 eV Ar+
Ti
I 20
40
60
Energy of sputtered particles E
80
(eV)
Figure 1.25. Energy distribution of sputtered atoms when bombarding a Ti substrate with Ar* ions with 1000 eV at normal incidence.
1.8. Experimental parameters in IAD thin film growth Hundreds of contributions concerning the deposition of thin films by using low kinetic energy ions exist in the literature. They mainly describe the experimental conditions used for the deposition of the films and their characteristics or properties. Several process parameters are very important for the control of the deposition conditions of thin films by any method based on the use of accelerated ion beams. In this section we are concerned with the experimental variables that are controlled in any ion assisted deposition of thin films. They have to be taken into account to make correlations between experimental preparation protocols and final properties of the films. Moreover, they also have to be considered to model the
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
41
mechanism of thin film IAD growth. Among the primary experimental IAD parameters, we identify
• • • • • • • •
The type of incident particles (atomic or molecular ions, neutrals, etc.) The kinetic energy or kinetic energy distribution of the incident particles The ion current density The total ion fluency The angle of incidence of the particle beam with respect to the surface of the target The ion to atom arrival ratio The temperature of the sample The ion energy momentum transfer
These experimental parameters are easy to correlate with the final properties of the films. Most commercially available ion sources permit a defined control of these primary process parameters that can be changed to properly modify the thin film characteristics in each deposition. The majority of these parameters (e.g., energy, type of ions, etc.) have a clear meaning and need not be discussed in detail here. In this section, only the concepts related to the transfer of energy and momentum to the growing film will be specifically discussed because they require a straightforward definition. In Chapter 3, some basic trends of the effect of all these process parameters on the characteristics of thin films grown under ion bombardment will be reviewed.
1.8.1. The ion to atom arrival ratio and the normalised energy concept In any ion beam assisted deposition process, a critical parameter is the ratio between the number of impinging ions per unit surface and unit time and the number of atoms effectively remaining on the film (due to sputtering this number will be smaller than the number of atoms actually arriving on the substrate surface). The ion fluency can be determined experimentally by measuring the beam current (i.e., /) at the sample position. The exact measurement of the beam current is not an easy task. For this purpose Faraday-cup probes are used (Gilmore 1995).
42
Low ENERGY ION ASSISTED FILM GROWTH
The flux of charged particles / impinging on the substrate surface, defined as the number of charged particles arriving at the sample surface per square cm and per second, can be estimated from the density current 7, at the sample position as I( particles/cm2s)
= 6.25xlOI2Ji(LiA/cm2
)
(1.43)
The energetic neutrals that may accompany the charged ions are a source of uncertainty for /. They can be very significant at high working pressures. In principle, it is expected for working pressures P higher than ~10"5 mbar the dependence / °c j . exp(ack,Pl/KBT)
(1.44)
where ac is the cross section for charge exchange, k, is a tooling factor, and I the distance from the ion source to the sample. According to Eq. (1.44), the higher the product PI, the higher the expected amount of neutrals in the ion beam. For example, in the case of 500 eV nitrogen bombardment with a Kaufman ion source for a total pressure of 2xl0"4 mbar and a distance of 26.4 cm from the ion source to the target, Van Vechten et al. (1990) estimated that 34% of the beam is neutralised. Another aspect to be taken into account when considering molecular ion bombardment is the effective number of atoms per ion. In this respect, for the same example mentioned above, it was estimated that their nitrogen beam was formed by 89% of N2+ species and 11% of N+ (Van Vechten et al., 1990). The flux of deposited atoms A is defined as the number of atoms remaining on the film per unit surface and unit time. An estimation of A depends on how these atoms are incorporated into the film (ion sputtering, sublimation by electron bombardment, vapour pressure of an volatile precursor, etc.). In principle, it can be obtained from the growing rate of the thin film. Although this magnitude can be deduced from the final thickness of the film and the total time used for its deposition, an "in situ" measurement of the actual thickness of the film at any time of the deposition process would be desirable. For this purpose, the use of a quartz crystal monitors, or techniques like ellipsometry can provide information about the thickness (Kasemo and Tornqvist, 1978; Fried et al., 2001).
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
43
It is also worth considering the residual gas flux Jr defined as the number of particles impinging on the sample surface per square cm and per second from the residual pressure Pr in the preparation chamber. Jr is related to Pr as Jr( particles/cm2s ) = 5.3xl02°
Pr(torr)
(1.45)
As mentioned before, from a practical point of view, it is interesting to know the ion to atom incorporated to the film ratio defined as I/A. Experimentally, it is found that many thin film characteristics depend on I/A. Figure 1.26 shows, as an example, the evolution of the width of the mixing zone for a carbon layer deposited on a silicon substrate by electron beam evaporation and simultaneously bombarded with 35 keV argon ions. The width of the intermixing zone, formed by a kind of SiC compound, progressively increases with the I/A ratio, for a constant energy of the impinging ions. The progressive accumulation of energy in that zone as the I/A ratio increases is the leading factor for such a tendency (Volz et al., 2000). In Chapter 3, a thorough description of the dependence on the I/A ratio of other thin film properties will be presented.
Figure 1.26. Width of the mixing zone, where recoil implantation of carbon into the silicon substrate and vice versa has been taking place. As a measure for the intermixed zone, it has been taken as the width of the region, where the carbon concentration is between 30 and 70 at.%. Reproduced from Volz et al. (2000) with permission.
Another general concept that can be found in many works on IAD thin films is the so-called normalised energy Enorm, defined as the average energy
44
Low ENERGY ION ASSISTED FILM GROWTH
deposited per atom incorporated in the film (Ji 1997). Emrm can be obtained from the expression E„on„=EiJl/A)
(1.46)
It takes into account the energetic contribution to the thin film growth of each individual ion impinging on the substrate. Thin film properties are highly dependent on this parameter as it measures the actual energy supplied to the growing film normalised by the number of atoms actually incorporated into the film.
1.8.2. Ion momentum transfer A typical magnitude of the accelerated ions that is very important for IAD is the socalled maximum ion momentum transfer PM. This concept takes into account the fact that, besides beam energy and ion fluency, the mass of the ions plays a definitive role in altering the properties of the bombarded targets. PM is defined according to: (Alvisi et al., 1999) PM=(l/Alj2M1TM
(1.47)
where, Mi is the atomic mass of the impinging ions and TM is the maximum energy transferred in a head-on collision as introduced in Eq. (1.12) Many thin film properties depend on the momentum transfer of the incoming ions. Thus, for example, preferential crystal plane orientation (Alvisi et al., 1999) or accumulation of residual stress in the film (Gerlach et al., 1998) have been found to be directly related with this parameter. Figure 1.27 shows, as an example, the dependence between the residual stress of TiN thin films prepared under bombardment with nitrogen or Argon ions of the same energy. A linear relationship is found between the residual stress and the maximum momentum transfer. Points in the figure for low momentum transfer were obtained for nitrogen bombardment, while those for high values of PM are obtained with Ar+ ions. It is assumed that the higher mass of Ar in respect to N is the main factor favouring the transfer of energy from the accelerated ions towards the film.
BASIC CONCEPTS ON THE INTERACTION OF LOW ENERGY ION BEAMS . . .
o-' (0 0CD
45
• nitrogen ions T argon ions
a.a
\
•
V.
•
£ 75 -3
T*.
1 ' ' '
0
100
200
•
•
•
300
'
•
-
•
•
400
•
•
1
'
500
'
'
•
600
maximum momentum transfer pmKL [(amu*eV)"2atom'1]
Figure 1.27. Residual stress of titanium nitride films prepared by nitrogen or argon-assisted deposition as a function of the maximum momentum transfer. Reproduced from Gerlach et al. (1998) with permission.
References Alvisi, M., Scaglione, S., Martelli, S., Rizzo, A. and Vasanelli, L., Thin Solid Films 354 (1999) 19. Coulter, C.A. and Parkin, D.M., J. Nucl. Mater. 88 (1980) 249. Eckstein, W., in Computer Simulation of Ion Solid Interactions, Springer Series in Material Science, vol 10 (Springer-Verlag, Berlin, Heidelberg 1991). Fried, M., Lohner, T. and Petrik, P., Ellipsometry Characterization of Thin Films, Ed. Nalwa H.S., Handbook of Surfaces and Interfaces of Materials vol. 4, p. 335, (Academic Press, San Diego 2001). Gerlach, J.W. et al., Surf. Coat. Technol. 103/104 (1998) 281. Gilmore, I.S. and Seah, M.P., Surf. Interface Anal. 23 (1995) 248. Hofer, W.O., in 'Angular, Energy and Mass Distribution of Sputtered Particles.' Topics in Applied Physics vol. 64: Sputtering by Ion Bombardment III, Editors R. Behrisch and K. Wittmaack, (Springer-Verlag, Berlin, 1991) pp.15-90. Hofsass, H., Feldermann, H., Merk, R., Sebastian, M. and Ronning, C , Appl. Phys. A 66 (1998) 153.
46
Low ENERGY ION ASSISTED FILM GROWTH
Ji, H., Was, G.S., Jones, J.W. and Moody, N.R., J. Appl. Phys. 81 (1997) 6754. Kasemo, B. and Tornqvist, E., Surf. Sci. 77 (1978) 209. Kittel, C , Introduction to Solid State Physics (Wiley, New York, 1976). Lennard, J.E. and Jones, I., Proc. R. Soc. London A106 (1924) 441, 463. Lindhard, J., Scharff, M. and Schiott, H.E., 'Range Concepts and Heavy Ion ranges' (Notes on Atomic Collisions II), in Mat. Fys. Medd. Dan. Vid. Selsk. 33 n°14, (1963) p. 3. Mitchell, T.E. et al., in Fundamental Aspects of Radiation Damage in Metals, M.T. Robinson and F.W. Young Jr. Eds. (US GPO Washington, D.C., 1976) vol.1 p. 73. Morse, P.M., Phys. Rev. 34 (1929) 57. Nastasi, M., Mayer, J. and Hirvonen, J.K., in Ion-Solid Interactions: Fundamentals and Applications, Ed. By D.R. Clarke, S. Sures and I.M. Ward FRS, Cambridge University Press 1996. Norgett, M.J., Robinson, M.T. and Torrens, I.M., Nucl. Eng. Des. 33 (1975) 50. Sigmund, P., in Sputtering by ion bombardment: Theoretical concepts in Sputtering by ion bombardment I: Physical sputtering of single-element solids, Editor R. Behrisch, Topics in Applied Physics vol. 47 (Springer-Verlag, Berlin, 1981) pp. 9-71. Sugiyama, H., Modifications of Lindhard-Scharff-Schiott formula for electronic stopping power, J. Phys. Soc. of Japan 50 (1981) 929, and Electronic Stopping power formula for intermediate energies, Radiation Effects 56 (1981) 205. Torrens, I.M., Interatomic potentials (Academic Press, New York, 1972). Vechten, D., van Hubler, G.K., Donovan, E.P. and Correl, F.D., /. Vac. Sci. Technol. A 8 (1990) 821. Volz, K., Kiuchi, M., Okumura, M. and Ensinger, W., Surf. Coat. Technol. 128/129 (2000) 274. Winterbon, K.B., Sigmund, P. and Sanders, J.B., Spatial Distribution of Energy Deposited by Atomic Particles in Elastic Collisions, in Mat. Fys. Medd. Dan. Vid. Selsk. 37 n°14 (1970). Ziegler, J.F., Biersack, J.P. and Littmark, U., The Stopping and Range of Ions in Solids (Pergamon Press, New York, 1985).
CHAPTER 2 ION ASSISTED METHODS OF PREPARATION OF THIN FILMS
In practice, the preparation of thin films by methods involving the assistance of the film growth with accelerated ion species can be carried out using several experimental approaches. Different methods can be considered depending on the procedure utilised for providing the material to be deposited and how the ions are generated, accelerated and directed towards the substrate. This chapter is dedicated to experimental aspects. We will first consider the methods of preparation of thin films consisting of the ion assistance of the growth of the thin film with ions provided by an independent source (section 2.1). These methods can be properly considered as "ion beam deposition procedures" (IBAD). In section 2.2, other IAD methods are presented where the kinetic energy and momentum supplied to the growing film are borne by the same particles that will be incorporated into the film. Owing to the increasingly higher technological importance of the Plasma Immersion Ion Implantation (PHI) procedures, a specific section is dedicated to briefly outlining the basic concepts involved in this type of deposition procedure. We present a brief description of the most typical ion sources utilised for IBAD deposition in section 2.4. The most important working parameters that have to be controlled for a reproducible preparation of the films were discussed in the previous chapter (cf., section 1.8).
2.1. Assistance of film growth with independent ion sources For this set of methods, the supply of the constituent material(s) of the film and its bombardment during growth are two independent, although simultaneous processes. A general scheme is presented in Figure 2.1. This figure shows that the deposited material arrives onto the substrate while, simultaneously, an ion beam impinges onto its surface. In this way it is possible to get separate control of the beam variables (i.e., energy of ions, current density, angle of incidence, type of ions, etc.) and those of the thin film deposition (basically the amount of material supplied to the substrate and the substrate temperature). Usually, these methods are referred to in the literature as Ion Beam Assisted Deposition (IBAD) (Itoh, 1989).
47
Low ENERGY ION ASSISTED FILM GROWTH
Figure 2.1. Schematic description of the growth of a thin film by an ion beam assisted procedure.
A first classification scheme of these methods arises from the consideration of the different ways of vaporising the material to be deposited and how it reaches the substrate surface. A non-exhaustive enumeration of the possibilities reported in literature is the following: Thermal or electron-beam evaporation Laser ablation Ion beam sputtering Volatile metal precursors The next sections present a description of the fundamentals, main parameters and the advantages of the different methodologies. However, before proceeding to this analysis, it would be interesting to summarise the order of the magnitude of the energy ranges of ion and/or particles involved in thin film deposition procedures and, in particular, IBAD methods. The scheme of Figure 2.2 shows the ranges of kinetic energy and equivalent temperature of incident atoms and ions involved in different thin film deposition methods depending on the procedure used for its generation. Evaporated atoms have an energy ranging between 10"2 to 1 eV. These values are very small when compared with the energies of ion beams used for assistance, usually ranging between some ten and some ten thousand eV.
49
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
In sputtering processes the atom energies are higher (typically from one to one hundred eV), while the ions may have energies ranging from ten to some hundreds eV. By pulsed laser deposition procedures the energy range of the evaporated species is also wider and can reach up to some hundreds eV. In direct ion beam deposition procedures, atoms are ionised and accelerated towards the substrate. The energy range in this case can be between some tens and some hundreds eV. Obviously, thin film properties will be very much dependent on the energy of the particles involved in its growth. Chapter 3 of this book sets out to discuss the type of effects produced by the ion bombardment and it will be shown that the energy is one of the most important parameters for an effective control of the final thin film properties. Energy (eV) 10* I
atom source
102
evaporation
I
evaporation
I
Iff1
1
10
102
1
II
I'
I1
104
105
106
103
equivalent temperature (K) I
i
sputtering
I
ion beam
deposition
ion source
t W # / % M i ion beam
l ^ t ^ m m ^ ^ m m i l
sputtering
10J I, 107
V//////////AV///////A
plasma plasma
W//W/////////////A on beam
llllllllllllllllllllllllllllllllllllllll i B ^ M * ^ ^ ^ * ^
ion beam laser deposition
Figure 2.2. Typical ranges of kinetic energy and equivalent temperatures of atoms and ions involved in different IBAD processes.
2.1.1. Evaporation and ion bombardment of the growing film By this procedure, the material to be deposited in the form of a thin film is supplied to the substrate by evaporation of a bulk precursor. Evaporation can be done by resistive heating of a crucible or, more efficiently, by electron beam bombardment of a bulk solid. Ion assistance of growth can be carried out with ion beams of a wide range of energies from low (E < 1 keV) to high-energy values (i.e., from E > 2 up to
50
Low ENERGY ION ASSISTED FILM GROWTH
40 or more keV). Usually, a given source can only work in a certain energy window and therefore different sources are required to cover a wide range of energies. Many low and high energy ion source designs, suitable for IBAD deposition, are now available (Ensinger, 1992). In section 2.4 we will present a brief description of some ion sources typically used for low energy ion assistance. High-energy ion sources are relatively more sophisticated and need additional requirements for operation. As for the low energy sources, the latest developments provide high current densities over a wide surface area. Depending on the model, they can also work with either rare or reactive gases. Low energy ion beams can be produced by means of ion sources such as those described in section 2.4. Generally, these are small experimental devices that can be implemented in deposition chambers of almost all research laboratories interested in these techniques of film deposition. Figure 2.3 shows a scheme of an experimental set-up where a solid is evaporated by electron bombardment and its atoms deposited onto a substrate while, simultaneously, the growing film is subjected to ion beam bombardment. Ion gun and electron beam evaporator are usually incorporated within the same chamber. Since the ion guns may operate at relatively low pressures (10"4 mbar or less), their working conditions are compatible with those of the electron evaporators (generally, to avoid arching in the electron gun the maximum working pressure in the chamber must be below 10"3 mbar). An experimental configuration such as that in Figure 2.3 permits a suitable control of the atom/ion arrival ratio as both evaporation rate and beam current can be monitored and controlled separately. It is also possible to use different types of inert (e.g., He, Ar, Ne, Xe) or reactive ion beams (e.g., N2+, 0 2 + , etc.). Other advantages of this procedure are connected to the high evaporation rate attainable with electron gun evaporators and the possibility of controlling the thin film properties by adjusting the evaporation rate, the ion beam current and energy and the temperature of the substrate during deposition. By a careful choice of the different parameters thin films can be prepared with a high rate and a good control of their properties. Thin films of different materials have been prepared by using this type of experimental configuration. Metals, oxides, nitrides and other compounds can be easily prepared as thin films in such an experimental set-up. For the preparation of oxide or nitride thin films, an advantage of this method is its flexibility in allowing a wide choice of the precursor material to be evaporated, which may consist of an oxide or nitride, as well as of a metal. In the latter case, the desired oxide or nitride
51
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
Substrate
lorn source1
Evaporator
™
Vacuum
Figure 23. Scheme of an IB AD experimental set-up combining electron beam evaporation and ion assistance with a beam provided by an- independent source.
compositions are obtained by assisting the growth with reactive ions, 0 2 + or N2+» depending on the type of material to be prepared (cf., section 3.8). Moreover, for a precise control of the stoichiometry of the film, a precise balance between the arrival rate of atoms and the ion current is required. Thus, under-stoichiometric thin films can be produced when the beam current is below the threshold value necessary to get the required oxide or nitride compound. Generally this threshold current is dependent on the beam energy because of the dependence of the preferential sputtering rate on the ion energy. The ion to atom arrival ratio (i.e., I/A) (cf., section 1.8.2) is another critical parameter for an effective control of the thin film composition. An example of this phenomenon is reported in Figure 2.4 showing the N/Ti ratio in a TiN thin film whose formation is assisted by bombardment with N2* ions of 5 keV under a working pressure of 2*10"3 Pa (Hubler et al., 1989). The amount of incorporated nitrogen increases with the ion to atom arrival ratio and
52
Low ENERGY ION ASSISTED FILM GROWTH
1
0.4
0.6
0.8
arrival ratio [N/Ti] Figure 2.4. Evolution of the N/Ti ratio in an IAD titanium nitride thin film as a function of the N/Ti arrival ratio ( • ) and comparison with the amount of nitrogen incorporated by adsorption of nitrogen while evaporating in the presence of this gas (A). Reproduced from Hubler (1989) with permission.
saturates for values greater than 0.4. In the plot, the amount of nitrogen incorporated by adsorption of N2 from the gas phase is also included. The solid line represents the calculated N/Ti atom ratio in the film when its growth is assisted with accelerated nitrogen ions. The theoretical curve results from the difference between the experimental values of the N/Ti ratio obtained by implantation and adsorption. It is apparent that at low I/A values the nitrogen in the film stems mainly from reaction of Ti with nitrogen gas, while at high I/A ratios implanted nitrogen coming from the accelerated ion beam is a majority in the film. 2.7.2. Laser ablation and ion bombardment of the growing film Another way of supplying the constituent atoms of the thin film to the substrate is by laser ablation of a precursor material. Laser beams may provide enough energy to a solid to vaporise its outermost layers. Pulsed lasers with wavelengths from 193 (ultraviolet) up to 1064 (infrared) nm have been used for evaporation of solid targets. Total energy of the pulse, with values ranging from some tenths to more than a thousand mJ per pulse, and the pulse width (typically some tenths of ns) and repetition rate (some tenths of Hz) are other parameters that are usually controlled during thin film deposition by laser ablation (Voevodin et al., 1996). When a suitable laser beam impinges on a point of the surface of a solid target, the high
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
53
energy which is supplied into it leads to its vaporisation. The evaporated material is then expelled from the irradiated surface in a rather collimated beam forming what is called a "plume". The plume is a small cone shaped region of space in the vicinity of the target where there are excited atomic, molecular and ion species that have been ejected from the solid and dragged out of it by a very intense thermal wave. The spatial distribution of species according to their energies can be determined by emission spectroscopy analysis of the plume (Kelly et al., 1992). Average kinetic energies and type of species can be modified by a proper control of the laser beam parameters. Clusters, single atoms or ions with different kinetic energies are present in the plume. Some typical examples of average kinetic energies of ion species are reported in Table 2.1 for graphite targets ablated under different conditions. Kinetic energy of a given species increases with the laser energy density, while the relative distribution of species of different types also changes with this parameter. From the point of view of IBAD processes a key issue is that the vaporised material can be partially ionised and that, therefore, it can be accelerated towards the substrate by applying a bias. Table 2.1. Main ion species and average kinetic energies produced by laser ablation of graphite under different conditions.
Main ion species
Average kinetic energy (eV)
Type of laser
Wavelength (nm)
C„+ C 15 +
3.3 4.5
Nd:YG
1064
C2+ C2+ C3+ C+
38 55 18 600
Nd:YG ExcimerKrF
532 248
ExcimerKrF
193
When the plume is oriented towards a suitable substrate, a thin film can be formed from the vaporised species of the target. Simultaneously to the deposition of these species, the growing film can be subjected to ion bombardment with an independent ion beam. A scheme of a typical experimental system combining laser ablation and ion beam assistance is shown in Figure 2.5. It consists of a laser beam that is focused inside a vacuum chamber where a target is located. The plume formed by the interaction of the laser beam and the target is directed towards a substrate that is simultaneously bombarded by an ion beam produced by an ion
54
Low ENERGY ION ASSISTED FILM GROWTH
source attached to the deposition chamber. With this experimental configuration, the adjustment of the atom/ion arrival ratio is quite straightforward. It just requires changing the power of the laser or its pulse cadence and/or the working conditions of the source (beam energy and current). In this way it is possible to get a precise control of the final characteristics of the film (texture, microstructure, etc.) under automatic operation. Other experimental parameters that can be adjusted are the temperature of the substrate and, if required, the value of a bias voltage to accelerate the ion species of the plume. This can be a cheaper, though less precise, alternative to the use of an external ion source.
Laser
Substrate
Tkrget Ion source (end-Hall)
Vacuum Figure 2.5. Scheme of an IB AD experimental set-up combining pulsed laser evaporation and ion assistance with a beam provided by an independent source.
At the present stage, the technique combining laser ablation with ion bombardment is mainly restricted to a laboratory scale because of the. relatively small deposition rates achieved with it. Another restriction of this technique is the small size area of the samples that can be obtained. This size is controlled by the
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
55
diameter of the plume, generally small at the relatively short distances separating the substrate from the target. The inhomogeneous lateral profile of species in the plume is another shortcoming of this technique that can limit the lateral homogeneity of the thin film properties. Experimental set-ups have been proposed combining substrate and target movements with optical dispersion of the laser beam with the aim of surmounting this drawback (Voevodin et al., 1996). Laser ablation is very well suited to vaporisation of refractive materials. In principle, owing to the high local temperatures produced at the point where the laser beam impinges on the solid, all its constituent elements are equally vaporised. Moreover, good vacuum can be preserved because the target is not extensively heated. These features make this technique very appropriate for deposition of thin films with a complex composition (mixed oxides, nitrides, etc.) and well-defined properties required for electronic, optical or optoelectronic applications. Typical examples of application of this technique are the preparation of mixed oxides, diamond and diamond like (DL) coatings, carbides, etc. Further developments and applications are expected for the future when some of the aforementioned problems of this technique are properly solved.
2.1.3. Dual ion beam deposition of thin films (DIBS) Another method of preparing a thin film assisted with ion beams is the so-called "dual ion beam sputtering" procedure (i.e., DIBS). By this method, the material to be deposited as a thin film is generated by ion beam sputtering of a target. One of the effects of the ion bombardment of solids is the removal of the topmost layers by sputtering. In Chapter 1, a description of the main features of this process was presented. In DIBS, the sputtered material is deposited on the surface of a substrate while, simultaneously, the growing film is bombarded with ions coming from a second ion source (Rossnagel, 1989). A scheme of a typical DIBS experimental setup is shown in Figure 2.6. Two Kaufmann-type or one Kaufmann and one end-Hall ion sources (cf., section 2.4.1) are generally utilised in this procedure. In more advanced developments, the use of more than two ion guns, for sputtering removal and assistance, has been proposed. Typically, the ion source used for material removal by sputtering operates at relatively high energies (of the order of 1 keV or higher) to maximise the sputtering yield (cf., section 1.7.1). The second gun, used to assist the film growth, operates at lower energies (i.e., some hundreds of eV)
56
Low ENERGY ION ASSISTED FILM GROWTH
to diminish the sputtering and to maximise the beneficial effects of the ion bombardment.
Vacuum Figure 2.6. Scheme of an IB AD experimental set-up combining ion sputtering and ion assistance with a beam provided by an independent source.
A critical geometrical restriction in a DIBS set-up refers to the relative orientation of the sputtered target and the substrate position. The geometrical arrangement of both elements has to take into account the spatial distribution of atoms sputtered from a target. The amount of material sputtered from a perpendicularly bombarded target roughly follows a cosine distribution function in respect to the target normal (cf., section 1.7.2). Therefore, to get a uniform distribution of the sputtered material on the substrate, the substrate surface must be located at a tangent position with respect to the almost spherical curve defined by such an angular cosine distribution function. For practical reasons, the angle between the substrate and target is, in most cases, around 45°. The substrate is usually rotated with respect to the perpendicular to their surface to get homogeneous thin films. To ensure thin film homogeneity, it is also critical that the ion beam profile used for assistance is homogenous. Another critical parameter of DIBS and
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
57
other ion beam techniques is the angle formed between the ion beam used for assistance and the substrate normal. Changing this angle provides a way of controlling the texture (i.e., preferential orientation of crystallographic planes, c.f., section 3.9) of the film. By DIBS the confinement of the two ion sources, target and substrate, in the same chamber may pose some practical restrictions for such alignments. However, modern DIBS instruments incorporate suitable elements to do such alignments in a very convenient way. For a proper control of the process parameters during DIBS deposition, the energy and ion density of the two ion sources must be very carefully adjusted. Thus, the amount of material supplied to the growing film depends, generally in a nonlinear way, on the ion energy and current of the high-energy gun. Meanwhile, the desired modifications in the thin film characteristics are produced by changing the value of similar parameters of the ion gun used for assistance. As in the previous procedures, the control of the stoichiometry for the formation of oxide or nitride compounds (i.e., metal/oxygen or metal/nitrogen ratios) is possible by carefully adjusting the oxygen or nitrogen ion current supplied by the second ion source. In that case, since the sputtering and deposition processes are carried out in the same chamber, a working requirement of the two ion sources is that they do not incorporate glow filaments to generate or to neutralise the ion beams (cf., section 2.4.3). Background pressure during deposition must be also very low (~ 10"6- 10"7 mbar) to ensure that the sources are working under the proper conditions of operation. The strict control of these and other process variables makes the quality of the films prepared by DIBS, in terms of planarity (i.e., low roughness), high densification or adherence to the substrate very good. With regards to these characteristics, DIBS is a very well suited technique for preparation of thin films for optical applications, high corrosion protection and, in general, when a very well controlled microstructure is required. The low roughness and high densification of thin films obtained by this method are important characteristics not only for single thin films, but also for the preparation of multilayers for different applications. However, a drawback of the technique is that, generally, the growth rate of the films is not high, since it depends on the relatively low supply of material generated by ion beam sputtering of a target. Since sputtering is usually less effective than other procedures such as thermal or electron beam evaporation, the growth rates achieved by DIBS are, in practice, lower than those obtained by these other methods. Control of the composition of a multicomponent thin film is not easy with this technique. In fact, since preferential sputtering effects may occur when a target
58
Low ENERGY ION ASSISTED FILM GROWTH
of complex composition (e.g., mixed oxides) is subjected to ion bombardment, the stoichiometry of the growing film (i.e., the ratio between the different metal elements present in the target) cannot be controlled as easily as by laser ablation. In any case, care must be taken to work under steady state conditions, once the composition of the outermost layers of the target has changed to compensate the preferential sputtering of some of their elements (cf., section 1.7.1).
2.1.4. Ion beam induced chemical vapour deposition (IBICVD) The previous three methods of deposition of thin films can be considered modifications of conventional physical vapour deposition (PVD) procedures where the growing film is being modified by the effect of an ion beam. In the three cases, the material to be deposited stems from solids that are either evaporated by heating or subjected to ion or laser irradiation. Another possible way of depositing a thin film is to use volatile precursors of the constituent element(s) of the film. Usually, these volatile precursors are chemical compounds with a relatively high vapour pressure at ambient temperature. Within this category one can find compounds such as metal halides, hydrides or metal-organic compounds. Typical chemical vapour deposition (CVD) procedures consist of the thermal decomposition of a volatile metallic precursor. By controlling the atmosphere and temperature of substrate during the decomposition, a metal, oxide or nitride thin film can be deposited, while volatile molecules of the other atoms constituents of the precursor (e.g., H 2 0, CO2, HC1, etc.) are removed from the reactor. Decomposition of volatile precursors can also be promoted by other methods such as plasma or laser activation (Lecours et al., 1993; Konuma, 1992). By using these activation procedures, the decomposition reactions and the deposition of the thin film may take place at lower temperatures. Ion beams can also induce CVD deposition. Owing to the fact that ion beams can be easily focused and/or rastered over a surface, this procedure was initially used for the deposition of thin metal wires for the repairing of integrated circuits (Shedd et al., 1986; Overwijk et al., 1993). A problem in the preparation of such wires is their contamination by carbon and other elements present in the volatile precursor. Silicon and silicon carbide have also been prepared by ion beam decomposition of appropriate precursors (Khan et al., 1999). Preparation of very pure metal oxide and/or nitride thin films is possible by IBICVD by using broad ion beam sources (Espinos et al., 1997). A scheme of an IBICVD set-up incorporating
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
59
one of these sources is shown in Figure 2.7. Main components in this experiment are the ion source and the dispenser of the volatile metal precursor. The latter must be constructed so that it provides a homogenous distribution of the precursor without making any shadow in the beam profile. Non-filament ion guns can supply low energy and broad 0 2 + or N2+ ion beams (c.f., section 2.4.3). ¥otottle precursor dispenser
Substrate
Ion source
Vncyiini Figure 2.7. Scheme of an IB AD experimental set-up combining chemical vapor deposition and ion assistance with a beam provided by an independent source.
The accelerated ions impinging on the substrate produce the decomposition of the molecules of the volatile precursor adsorbed on the surface of thefilmwhile, simultaneously, assisting its growth process. Temperature is an independent working parameter that can be adjusted according to requirements, although at high values the growth rate may decrease since the sticking coefficient of the precursor molecules usually decreases. In general, the oxide or nitride thin films obtained are very compact and free from contaminating elements, which are removed from the growing layer in the form of simple gas molecules. An advantage of this method in comparison with the previously described IBAD procedures is its relatively lower
60
Low ENERGY ION ASSISTED FILM GROWTH
cost. By this procedure, only one ion source is required, while two sources are used in DIBS or one ion source and electron beam evaporation or laser ablation devices in the two other methods described previously. However, a disadvantage of this method is that contamination of the ion source may occur if the pressure of the volatile precursor increases above a certain limit. Since the ion sources are supplied with either 0 2 or N2, another restriction of the method is that they cannot use incandescent filaments for operation. Typical ion beam energies utilised for the deposition of thin films by IBICVD are above 300 eV. A careful balance between the ion current at the substrate position and the partial pressure of the precursor in the chamber is necessary to prevent sputtering prevailing over deposition with the result that the film is not formed.
2.2. Ion assisted deposition of thin films without independent ion sources Experimentally, kinetic energy and momentum can be supplied to a growing film by using methods that do not require the use of independent ion sources, but rather the impingement of accelerated species produced by other procedures. In this section we will review some of these methods. Several possibilities will be considered depending on whether the film is subjected to bombardment with inert or reactive species which are firstly produced in a plasma and then accelerated towards the substrate by an electrical field, or the direct ion beam deposition techniques whereby the metal atoms or clusters to be deposited are themselves ionised and accelerated before impinging on the substrate surface. Within the first type of procedures we will consider the so called "ion plating" and "ionised magnetron sputtering" procedures. From the second type we will refer to the "ionised cluster beam", "filtered arc" and other direct and mass selected "ion beam deposition" procedures. Complex situations can also be found where both plasma species and ionised atoms arrive simultaneously at the substrate surface. By these alternatives biasing of the substrate, plasma, independent ionisation and acceleration of metal species, etc., are conveniently combined. Further details of different possibilities can be found in more specialised works (Zhurin et al., 2000).
2.2.1 Ion plating The term "ion plating" has been used as a general description of those experimental
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
61
procedures of deposition, where a growing thin film is subjected to a flux of highenergetic particles. In this respect the methods of deposition discussed previously in section 2.1, where the growth of the thin film is assisted with an accelerated ion beam supplied by an independent ion source, are sometimes considered within this term. Here, we will restrict the "ion plating" concept to those methods where the deposition of the film is carried out within an inert or partially reactive gas discharge forming a plasma. The substrate is embedded within the plasma and eventually biased to accelerate its positively charged particles (Mattox, 2000). Evaporation of the deposition material can be done either resistively or, as is more common, by electron beam evaporation. When the material to be deposited is resistively evaporated the vacuum restrictions are not so critical and relatively highpressure plasma can be formed by DC biasing or other alternative methods. The base pressure of the system constitutes an important restriction when using electron beam evaporation. Typically, electron beam evaporation sources cannot work at pressures higher than 5x10"" mbar. Therefore, this is the maximum gas pressure that can be attained in the chamber in the vicinity of the electron evaporator source to generate the plasma. Plasma excitation of a gas can be achieved at such low pressures by the so-called Electron Cyclotron Resonance (ECR) procedures (Holber, 1989). Sources relying on this concept usually work in a down-stream configuration and are attached as an external device to the deposition system. However, since for ion plating it is convenient that the plasma fills the whole deposition chamber, other alternative procedures of generating a gas discharge are also used. Figure 2.8 shows the scheme of a commercially available deposition system where the plasma is formed through a low voltage discharge in a source attached to the main chamber (Pulker, 1999). Typically, an Ar plasma is formed between the filament of the source acting as a cathode and the electron beam crucible acting as an anode. In this way, the excited gas zone may fill the whole chamber. The substrate holder is electrically floating. Under these conditions, a plasma sheath (i.e., a thin plasma region surrounding the substrate that is depleted in electrons in respect to the bulk of the plasma) is formed at the surface of the substrate and the substrate potential has a negative value of about 15-20 V with respect to the plasma volume. This voltage, plus the effect of the repulsive field produced by the anode, acts on the positively charged species of the plasma. In this way, the charged particles become accelerated towards the substrate and impinge on it with a given kinetic energy. When reactive depositions are intended (e.g. deposition of oxide materials) some oxygen or another gas can be added to the vessel.
62
Low ENERGY ION ASSISTED FILM GROWTH
Figure 2.8. Scheme of an ion-plating experimental set-up combining electron beam evaporation and generation of a plasma.
An interesting effect that appears because of the immersion of the different deposition elements within a plasma is the partial ionisation of the evaporated metal atoms that are forced to traverse the plasma zone before arriving at the substrate. Therefore, not only neutral species of the material to be deposited may reach this latter, but also ion species that are accelerated by the electrical field developed between the source and the substrate itself. Besides that, accelerated species from the plasma (i.e. Ar* and/or 02*) impinge on the film during its growth and contribute to its densification and to the other beneficial effects that make the IAD methods so interesting (cf., Chapter 3). A scheme of some of the physical processes involved during the deposition is presented in Figure 2.9. According to this scheme, electrons, Ar atoms and ions and neutral and ionised metal atoms arrive at the substrate surface. The metal ions are produced during the passage of the vaporised
ION AS1STED METHODS OF PREPARATION OF THIN FILMS
63
EquipotentiaJ lines
cruciable
Figure 2.9. Scheme of the processes occurring by the deposition of a thinfilmby ion plating. M: evaporated metal, S: substrate material, Ar: argon atoms or ions.
metal atoms trough the plasma. Two mechanisms have been described to account for this ionisation: electron-atom collisions (i.e., e" + MO -> M+ + 2e") or the socalled Penning ionisation (i.e., MO + G* ™> M+ + G° + e\ where G* refers to Ar or any other excited atom species from the plasma). Much research has been carried out to estimate the degree of ionisation of the metal atoms before arriving at the substrate. The reported values range between 30 and 0.03 %. The most realistic results seem to be of the order of 1% or less. Properties of the thin film (hardness, adhesion, density, etc.) are very much dependent on the degree of ionisation. The ionisation percentage can be changed experimentally by modifying process parameters such as the type of plasma gas or the gas pressure and electrical field strength (Ahmed, 1987). The number of ionised metal atoms plays a vital role in reactive ion plating processes, especially in low temperature reactions aiming at the synthesis of oxide or nitride thin films. Another characteristic of an ion plating
64
Low ENERGY ION ASSISTED FILM GROWTH
process that is very interesting for the control of the thin film properties is the energy distribution of the Ar+ ions in the plasma. Usually, the distribution profile is characterised by an exponential decay of the percentage of ions as a function of their energy. The maximum energy corresponds to the discharge voltage, while the average energy ranges around 10% of that maximum energy value, depending on experimental parameters such as gas pressure, electrical field, etc. In comparison with the aforementioned procedures based on the use of independent ion sources (cf., section 2.1), the energy, momentum and other characteristics of the species involved in the deposition are less controlled in the ion plating methods. This fact stems from the complex energy distribution function of the ions and because other process parameters such as the ion current (i.e., number of ion species impinging on the target) cannot be directly pre-established. By contrast, ion-plating methods usually provide a higher film growth rate and are less size restricted than the methods using independent ion sources where the beam diameter determines the sample size. Here, the plasma volume usually occupies the whole chamber and therefore larger samples can be covered with the species coming from the crucible. Another interesting feature of the ion plating procedures is that deposition may also occur in out of line directions. In fact, according to the schema in Figure 2.9, additional deposition lines are created through the electrical field lines leading to the back or other zones of the substrate. However, we must bear in mind that homogeneity is not preserved in these out of line sample regions.
2.2.2. Ionised magnetron sputtering (IMS) Magnetron sputtering is a widely used method of deposition of thin films that is utilised for a large variety of applications at laboratory and industrial levels (Kelly et al., 2000). High deposition rates, easy scaling, possibility of depositing metals as well as insulator materials and good quality of the deposited films are some of the advantages that have made this method so popular in many fields of thin film research and technology. A magnetron-sputtering device consists of a target (or cathode) plate that is bombarded by energetic ions generated in a glow discharge plasma, situated in front of the target. The bombardment process causes the removal (i.e., "sputtering") of target atoms, which may then condense on a substrate as a thin film. Most species sputtered from the target are neutral atoms. For certain applications, such as, for
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
65
example, the semiconductor industry where high aspect ratio features have to be homogeneously covered (the aspect ratio is considered as the depth/width ratio of the feature) magnetron sputtering does not provide the desired finishing (cf., section 3.3.3). To overcome this problem and to deposit metals into trenches and tracks of high aspect ratio for interconnects in electronics, the classical magnetron sputtering procedure was modified into the so-called "ionised magnetron sputtering" (IMS) method, also known as "ionised metal physical vapour deposition" (IMPVD) (Rossnagel et al., 1993). In IMS, physical sputtering, typically from a magnetron cathode, produces a flux of metal atoms in the direction of the substrate. A secondary plasma, typically inductively coupled (ICP), is produced between the target and the substrate by a radio frequency (RF) driven antenna. A scheme of the experimental set up is shown in Figure 2.10. It consists of a sputtering source, a two turns coil of a large size to generate the plasma and a holder with the substrate placed on it. The holder can be either biased or located behind a series of grids. Both configurations aim to accelerate the ionised metal atoms before they impinge onto the substrate. In the experimental configuration of Figure 2.10, the plasma is sustained by the inert gas (Ar, Ne) used for sputtering. Since the pressure is in the range of some tens of mbar, the sputtered atoms are slowed down and ionised before they reach the substrate. Preferential ionisation of the metal atoms by the free electrons of the plasma occurs because they have an ionisation potential (IP) smaller than that of the inert gas (e.g., IP of Al is 5.98, while that of Ar is 15.75 eV). Depending on the working conditions, ionisation fractions may vary from 10 to 90% of all metal atoms. For a constant pressure of the inert gas, the ionisation fraction increases with the RF power and decreases with magnetron power. After collision with the atoms, the plasma electrons lose a significant part of their energy producing the cooling and, eventually, the complete quenching of the plasma. Magnetron sputtering is considered a physical vapour deposition procedure. The IMS method constitutes a very simple and efficient modification of this well tried method that provides an efficient way of assisting the growth of a thin film with ions formed during the same deposition process. A clear advantage of this procedure is that it only requires a simple modification of a conventional magnetron deposition system consisting of incorporating a RF field within the deposition chamber. A disadvantage in relation to those IBAD methods that incorporate an independent ion source for assistance of the growth process (cf, section 2.1) is that the energy of the ion species and the effective current at the sample position are not well defined and their control is not straightforward. Another problem of the method
66
Low ENERGY ION ASSISTED FILM GROWTH
is contamination. In this respect, a careful adjustment of the working conditions is necessary to avoid the sputtering from the material of the immersed coils.
I F Colli
Vacuum Figure 2.10. Scheme of an IAD experimental set-up combining magnetron sputtering, plasma ionisation and acceleration grids.
2.2,3 Filtered vacuum arc deposition (FVAD) In the aforementioned methods, ion plating or ionised magnetron sputtering, the metal ions reaching the substrate are produced by interaction of vaporised metal atoms with Ar plasma. The vacuum arc deposition method can be considered by itself as another IAD procedure since it involves the accelerated ion species that has been produced by ignition of a vacuum arc. A typical vacuum arc is produced when a high current arc (e.g. hundred Amps) flows between an anode and a cathode after applying a voltage difference between them (typically some tens of volts). This high current intensity first melts and then evaporates the material from the cathode surface. Within the discharge region, a plasma arc is formed with the evaporated
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
67
material. The cathode discharge spot can be regarded as a micron-sized source of a neutral plasma consisting of electrons, metal atoms, metal ions and even micronsized droplets of the cathode material. Table 2.2 lists some typical values of current density, electron temperature, pressure, ion energies, etc. found in an arc discharge (Martin et al. 1992). In terms of thin film deposition, the most interesting features of a vacuum arc are the charge state of the ions (the arc produces multi-charged species), the ion fraction and the ion energy distribution. The actual values of these parameters are strongly dependent on the type of evaporated material and the figures quoted in Table 2.2 should only be taken as an orientation. Table 2.2. Standard parameters of a cathode spot discharge
Parameter Current density Electron density Electron temperature Crater size Ion energy Ion fraction
Range 107-1010Am"2 5-1020 m 3 (Cu) 3-6 eV (Cu) 6-9 eV (Al) 1-20 u.m 25-75 eV 0.1-10
The presence in a vacuum arc of metal atoms and droplets besides ion species makes this technique unsuitable for the preparation of thin films with strict specifications. In particular, particle contamination due to the deposition of the vaporised metal droplets may be deleterious for many thin film properties. To avoid such contamination, several experimental approaches have been developed with the aim of removing the neutrals and micro particles from the vaporised source, while leaving the ionised species unaffected (Martin et al., 1992). The filtered vacuum arc deposition (FVAD) constitutes a very effective procedure for this purpose. A typical experimental set-up is shown in Figure 2.11. It consists of a vacuum arc source composed of a cathode, an anode and a trigger electrode to initiate the arc. A magnetic filter serves to remove the neutrals and the droplets, while only the ionised species can traverse it and reach the substrate holder. Although there are several designs of filters based on different concepts (Sanders et al., 2000), the most popular one is the magnetic filter based on the application of a toroidal magnetic field parallel to the wall of a torous. In a system of this type, the magnetic field forces the charged particles to describe a circular trajectory preventing their collisions with the torous walls. In the presence of a magnetic field, the electrons spiral around the
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magnetic field lines until they suffer a collision with another particle. If the magnetic field is bent, the electrons follow the curvature. The positive ions are forced to follow the magnetic field lines due to the electric fields developed between electrons and ions. The plasma stays macroscopically neutral. Plasma transport is therefore the result of a combined effect of magnetic and electric fields. Contrary to the charged species, macro particles move along straight trajectories and cannot follow the curvature of the plasma duct.
Figure 2.11. Scheme of an IAD experimental set-up combining vacuum arc evaporation and mass filtering.
The geometrical condition that must be fulfilled for an effective transport of a low-density plasma stream along the toroidal field of a magnetic plasma duct is given by: r/u>l/v0
where
u=(Mcv02)/(ZeRH)
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
69
where u is the velocity of the centrifugal ion drift in the field; M, Z and v0 are the mass, charge and longitudinal velocity of the ions; R is the radius of curvature of the magnetic lines of force; / the length of the toroidal field and r the radius of the plasma duct. Typical ion currents provided with a set-up of this type are in the order of 400 mA, or even higher if a cross electric field is applied to the system. A clear advantage of this method is the high deposition rates attainable with it. A difficulty is that the control of some deposition variables (e.g. ion energy) is not always possible. However, an experimental configuration, such as that in Figure 2.11, offers the possibility of further accelerating the ion species by applying a bias voltage to the substrate. Another alternative is to use an additional ion source to assist the deposition of the growing film by bombardment with the ions supplied by that independent source. This alternative opens the possibility of depositing oxide and nitride materials by assisting their growth with either 02 + or N2+ ions supplied by the ion source. Sometimes, the high ion currents impinging on the substrate in a FVAD set-up may locally heat the growing sample to a high temperature. Besides that, the exact value of the local temperature cannot be determined properly, such heating may have deleterious consequences on the characteristics of the films and may lead to their delaminating or to other undesirable effects. In any case, a precise control of the ion current and other deposition conditions is required for each type of material.
2.2.4 Ionised cluster beam (ICB) In the FAD procedure, the metal ions may have different ion charges and are characterised by a wide energy distribution function. So, it is not always possible to establish a precise correlation between the characteristics of the film and the process parameters. The energies of the ions are in any case small, usually below 100 eV (cf., Table 1.2), except when a biased voltage is applied to the substrate to accelerate charged metal ions before reaching the growing sample. More precise control of the growth process can be achieved with the socalled "ionised cluster beam" (ICB) deposition procedure (Takagi et al., 1975), a review of which has recently been published (Yamada et al., 2001). By this method, metal clusters with about a thousand atoms are ionised, typically with a single charge, and accelerated under the action of an electrical field to energies of some keV. Averaging this energy per each single atom of the cluster yields a small energy
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Low ENERGY ION ASSISTED FILM GROWTH
per atom, of the order of some eV. Owing to the low energy carried by each atom, the damage produced on the substrate surface is small. Therefore, this method is especially suited to preparing very dense and smooth thin films. Epitaxial growth is also a typical application of this technique. Figure 2.12 shows as an example a molecular dynamic simulation of the interaction of an accelerated 2000 Cu atoms cluster with a Cu(100) surface (Moseler et al, 20(D). The acceleration energy of the cluster is 10 keV and the average energy per atom about 5 eV. During the initial stages of the interaction, the cluster impinges with a very high temperature and pressure over the surface and spreads laterally. Although the initial damage produced at the substrate extends over ten atom layers, relocation and readjustment processes serve to release the accumulated energy. In the end, a well-ordered layer with well-defined interface and little intermixing with the substrate is formed.
Figure 2.12. MD simulation of the interaction of an ionised cluster with a surface. Reproduced from Moseler et al. (2000) with permission.
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
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A typical experimental set-up for ICBD is shown in Figure 2.13. It consists of an evaporation and cluster formation zone, an ionisation area and an accelerated stage. Initially a metal vapour is ejected into a high vacuum region through a small aperture of a special crucible where metal is vaporised by resistive or electron bombardment heating. The clusters are formed from the metal vapour through adiabatic expansion and atom collision (i.e., cooling wall stage). Then, they become ionised by collision with electrons that are emitted from a filament coil located coaxially in front of the crucible at a negative voltage Va with respect to it (i.e., stage). A negative potential V, is also applied to the substrate to attract the ionised clusters. Under these conditions they reach the sample surface with a kinetic energy of e(Va+Vi). Adjustment of the evaporation and clustering processes to produce a very narrow distribution of cluster sizes is critical for precise control of the thin film characteristics.
Vacuo in Figure 2.13. Experimental set-up for ICBD.
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The preparation of oxide or nitride thin films by reactive ion cluster beam deposition (RICBD) is also possible by supplying oxygen or nitrogen at the sample position. As for the metal thin films, a typical feature of oxide and nitride thin films prepared by ICB deposition is their smoothness and the possibility of preparing epitaxial thin films grown with a certain preferential orientation of some crystallographic planes (cf., Section 3.9) (Cho et al., 1999). Recently, the experience gained in the preparation of fullerenes and similar carbon clusters has been used to prepare carbon films by ICBD. In this case, fullerenes are directly evaporated and then ionised before impinging on the substrate surface (Maiken et al., 1995). The high vacuum requirements of this technique and the high cost of the overall set-up have limited its use for the preparation of very well defined layer or multilayer structures. An advantage of the technique is that it is compatible with the use of "in-situ" diagnostic methods based on the use of electrons and requiring very strict vacuum conditions. Thus, it is very common to find ICB systems where the growing process of the film can be followed "in-situ" by means of reflected highenergy electron diffraction (RHEED).
2.2.5. Mass selected Ion Beam deposition (MSIBD) This method, also known as Ion beam deposition (IBD), consists of producing ion beams of a given element with very small energy dispersion. The chosen element(s) become incorporated into the substrate where they form a layer (Marton, 1994). This method is usually chosen when well-defined films with good adhesion and controlled characteristics have to be prepared. The most important parameter by this method is the ion energy. Usually it ranges around 100 eV, a value for which sputtering is not yet significant though the sticking probability of the ion species on the surface is high (of the order of the unity, meaning that all the impinging ions will remain on the substrate). Probably the most appreciated effect of the IBD method is that the ions within that energy range can penetrate below the first layer(s) of the substrate where they become incorporated. This feature constitutes the basis of the so-called sub-plantation effect that will be dealt with in detail in section 5.12.3. The growth of the film beneath the surface confers to it interesting characteristics in terms of good adhesion, the possibility of formation of new phases, the production of very smooth surfaces, etc. The reduction of the sputtering yield to a minimum is very important for the achievement of such goals, it also being important that the growth process may occur at sub-surface regions.
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
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When an accelerated ion approaches a surface there are a series of processes that may occur before and during the incorporation of the element within the layer. Among these processes we can mention affecting the charge exchange, the ion impact and the thin film growth. When an ion species approaches a surface it becomes neutralised at a distance of a few Angstroms from it. Neutralisation is straightforward when the substrate is a ground metal, but it may represent some problems when it is an insulator where charge might be accumulated and eventually produce some arching. As an effect of ion impact the impinging species may either become adsorbed onto the substrate surface, penetrate beneath the surface or be back scattered from the substrate. For the ion energy range used by IBD, back scattering uses to be negligible, while penetration only takes place if the ion energy is above a certain penetration threshold value typical of each material and ion species. For the energy values used in this technique, virtually no collision cascade (cf., section 1.6.1) is produced within the target, but some knock-on and related binary processes that lead to other phenomena such as displacement of atoms and formation of lattice defects. Such processes only occur if the ion species have the minimum energy necessary to produce such displacements once they have penetrated below the surface. The maximum energy transferred in a head-on collision and the minimum kinetic energy of the impinging species necessary to produce atom displacements (i.e., displacement threshold energy, Eth) can be estimated according to Eqns. (1.12) and (1.13) respectively. These equations can be used to estimate the optimum energy values of the ions for the production of thin films by ion incorporation below the surface. If the ions have a kinetic energy Ek<Eth, they can only be stopped within the target lattice by incorporation in interstitial sites. Ions of this nature are optimal for the deposition of high quality films because, even if they penetrate the substrate, they do not create Frenkel pairs involving the formation of atom vacant positions in the lattice. However, for practical reasons, the actual energy of the ions is usually much higher than £,A. Once the ion species have penetrated within the surface, several processes can be involved in the formation of a thin film. First is diffusion towards the surface. Thus, the excess of atoms incorporated in regions near the surface by ion bombardment may diffuse away to the surface to compensate the excess of surface energy. Actually, for the synthesis of many thin films, this process is not desirable because thin films grown on the top of the substrate surface are less dense than those grown in sub-surface regions. Growth of metastable phases can also be
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Low ENERGY ION ASSISTED FILM GROWTH
hindered if diffusivity is high. A clear example of this is the growth of diamond films for which the substrate cannot be heated during deposition to avoid diffusion of carbon atoms to the surface where they would aggregate as graphite (cf., section 5.23). At present, IBD procedures are still expensive and are mainly used for research purposes to get very good quality model thin films and/or to study basic processes of thin film formation. Nevertheless, some applications have also been reported for the synthesis of GMR reading heads (cf., section 4.7.2). Carbon, metal, semiconductor (Ge and Si) and some compound (SiC, III-V compounds, BN) thin films have been prepared with this method. The experimental set-up required for this procedure consists of the following elements: an ion source, a mass selector element, an ion deflector section and an ion decelerator. These parts are schematically represented in Figure 2.14. The whole system works under UHV conditions. The ions, once they have been produced by the ion source, are massselected (a classic compilation of beam sources for heavy ion production is that of Freeman et al, 1977; other designs for ion sources have been reported by Ensinger, Extractor
Ion sottFce
Mass selector Flight tube
Decelerator Substrate
1®'
II"
it-
it-
IfT
II
Pressure (Torr) Figure 2.14. Scheme of an IBD set-up.
1992). The mass selection section, consisting of a magnetic mass analyser, is required to produce a pure beam of ions for deposition. The need for such a set-up is obvious when metal ions such as C+, Si+S etc., are the selected species for deposition
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
75
and have to be separated from other gas residual ions. Before being mass-selected, the ion beam extracted from the source is accelerated from the source potential (i.e., Us) to the flight tube potential, usually to a few thousand volts. It is therefore necessary to decelerate the beam of ions before it impinges the target. This is done in a combined system consisting of a deflector and a decelerator section. In the deflector section, the ions are separated from any neutralised species present in the beam. Then, deceleration up to ground potential is carried out, either with a magnetic or an electrostatic lens system. The effective kinetic energy of the ions when they reach the target surface is given by the source potential, i.e., Ek-Us/e~. Either the extraction-accelerator or the decelerator parts have to be constructed so that they produce rather mono-energetic ion beams free from any other impurities, neutrals, etc. For some experiments involving the simultaneous deposition of two elements, multiple-source systems have been constructed. They usually consist of two independent lines, in which case the orientation of the substrate towards the lines and the adjustment of the deposition conditions are critical for a precise control of the stoichiometry of the film.
2.3. Plasma immersion ion implantation The plasma immersion ion implantation (PHI) technique was developed in the mid eighties for the surface modification of different materials such as metals, ceramics, etc. Since the thickness of the surface layer that has to be modified by this technique is of the order of several tenths of a nanometer, the energy of the ions must be relatively high. In this respect, this technique presents some similarities the ion implantation methods, where a beam of highly energetic ions (N2+, 0 2 + , Ar+, metal ions, etc.) of some tenths or even hundreds of keV is extracted from an ion accelerator to be implanted within the external layers of a target (Riviere, 1992). However, while the conventional ion implantation is a line of sight method, the PHI technique avoids such restriction and permits implantation and deposition by using a much simpler and less expensive deposition system. In this book dealing with IAD methods, we include a discussion of the PHI procedure because it can be modified to produce not only ion implantation but also ion-assisted deposition of thin films.
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2.3.1 Plasma immersion ion implantation (PHI) The PHI technique developed by Conrad and co-workers (Conrad et al., 1987) is not a line of sight process and does not require the use of expensive ion accelerators. In principle, no size or shape restrictions exist in PHI. The target piece is placed within a plasma and is pulse biased to a high negative potential (i.e., several tenths of keV) relative to the chamber walls. Plasma ions become accelerated towards the target where they become implanted. High ion implantation doses can be obtained by this procedure. Since the total time of application of the pulses is small compared with the total time of operation, extra heating of the sample is avoided. When a large negative potential pulse is applied to a target piece placed within a plasma, a plasma sheath develops around the target. Acceleration of ion species across the plasma sheath surrounding the target means the line-of-sight restrictions of conventional ion implantation are avoided. Three different stages have been considered in the evolution of the plasma sheath within a pulse (Le Coeur et al., 2000). Figure 2.15 shows a scheme of the evolution of the plasma sheath in relation to the magnitude of the ion current arriving at the target as a function of time. The voltage profile through a pulse is also reported in the figure. The plot shows that, initially, there is a drastic and sharp increase of the current followed by a nearly constant value until the pulse vanishes. During the initial stage of formation of the sheath, electrons are repelled from the target, leaving behind an unbalanced number of positive ions in an ion matrix sheath. These ions become accelerated towards the substrate. The energy distribution of the ions when they reach the surface of the target depends on their initial position within the plasma sheath. As the target collects the ions, new ions are extracted from the plasma and, therefore, the length of the plasma sheath increases. On a longer time scale within the pulse, a steady state is reached (the so-called Child-Langmuir state) and the sheath remains static until the voltage is removed. An interesting magnitude for the operation of the PHI technique is the maximum thickness that the sheath may reach within a pulse. In a first approximation, it can be estimated that the maximum sheath thickness is given by:
g = XJeV0/KTef
(2.1)
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
77
where V0 is the potential of the substrate during the pulse, Te is the electron temperature of the plasma and Xye the electron Debye length. As an example, for V0 = 100 KV the sheath thickness may reach a value of 40 cm. The expected value of the sheath thickness has to be considered for properly scaling the size of the experimental reactor that has to be scaled to contain a plasma volume greater than the sheath region.
Formation of matrix sheath
Sheath expansion
Steady state (Ghild-Langmuir)
Figure 2.15. Evolution of the current and voltage through a pulse in a Pill experiment. The formation process of the plasma sheath is schematically represented.
For good control of the implantation process it is also necessary to know the energy distribution function of the ions. Implantation with almost monoenergetic ions is possible by adequate control of the pulse variables. A condition for monoenergetic implantation is that the transit of ions through the sheath is collisionless. This implies that the mean free path of ions through the plasma is larger than the plasma sheath length. To achieve these conditions, low-pressure plasmas are preferable. The use of ECR plasmas, rather than hot filament plasmas, is recommended for such purposes, especially in the case of the production of reactive plasmas.
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2.3.2. Plasma immersion ion deposition (PUD) Previous considerations apply to the implantation of ions from a plasma gas. The final result is the formation of an implanted layer (of an oxide, nitride, etc.) with an average composition that depends on the chemical characteristics of the plasma. In this sense this technique has to be considered as an implantation rather than a deposition technique. The PHI concept can also be applied by positioning the target within a metal plasma. Under these conditions, implantation of accelerated metal ions takes place during the time of application of the pulse to the target, while deposition of metal species takes place when the sample is not being pulse biased. The interaction of accelerated metal ions during the pulse period with the previously deposited metal layers can be considered as an ion assisted process where phenomena such as energy and ion momentum transfer will contribute to modify the physico-chemical characteristics of the layer. A wide range of material modifications can be obtained by adjusting the implantation/deposition duty cycle and the magnitude of the substrate bias voltage. Moreover, if a reactive gas is introduced into the chamber (e.g. 0 2 , N2), a film of a new compound can be produced. Metal plasmas can be produced very efficiently in vacuum arc set-ups combined with PHI devices. Figure 2.16 shows a scheme of an experimental facility that combines a filter vacuum arc apparatus to generate a metal plasma and a pulsed substrate to induce the implantation of ion species from the plasma (Anders, 1997). Pulse characteristics of the vacuum arc source and the magnitude of the substrate bias are also compared in the figure. It is apparent that each arc pulse can be synchronised with the bias pulses of the substrate. Tuning the two types of pulses can change the overall efficiency of the process. In the pulse scheme of Figure 2.16 the two types of pulses have been adjusted so that biasing the substrate only occurs in the presence of metal plasma. In this way, pure ion implantation takes place without any film deposition. There are other variants for the adjustment of the pulse arc that lead to a large diversity of processes and, conversely, thin film properties. If implantation and deposition phases are alternated during the implantation phase a freshly deposited film is bombarded with energetic ions, thus leading to the formation of an intermixed layer that would be responsible for a superior adhesion of these films. When working with vacuum arc plasma, the sheath formation mechanism proceeds in a similar way to the "conventional" gas plasma version of the PHI technique. However, some peculiarities deserve a specific comment. In these
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
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plasmas, multiple charged Ion species are very abundant and the overall degree of ionisation is very high, sometimes approaching total ionisation. Under such conditions the ions Impinging on the target will have different energies according to their charge and the sheath thickness Is smaller than In conventional plasma.
High Voltage
Sequences of f tie high visltage pulsing
Vacuum
Figure 2.16. Scheme of an IAD system Integrating a FVA and a pulsed substrate to induce Ion Immersion processes (I.e., PHD).
PHI and deposition has also been attempted by combining magnetron sputtering with pulse biasing the substrate or by evaporation of a metal within a plasma surrounding the substrate (Brown et al., 1999). As an example of the possibilities of this latter technique. Figure 2.17 shows an SEM micrograph of an A320:r2S102 mullite thin film deposited on SiC. The film is prepared by multiple FVA from two vacuum arc sources of Si and Al and by adding some oxygen in the deposition chamber. The high compactness and planarity of that film is apparent in the figure. Perfect control of the stoichiometry of the deposit and an enhanced thin film adhesion are some other features of this type of thin film preparation that was achieved by the synthesis of this mullite material.
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Figure 2.17. TEM micrograph of a cross section of a AR> r 2SiO- mullitc thin film deposited on SiC by PIID. Reproduced from Brown et al. (1999) with permission.
2 A Broad beam ion sources The development of broad beam ion sources during the seventies can be considered as a critical turning point in the evolution of the ion beam assisted procedures of depositing thin films. The straightforward operations of these sources and their relatively low price have favoured the expansion of their use for the growth of IB AD thin films. In this section, we will review their general principles of operation and comment on the most utilised designs for thin film preparation. The ion beam sources used for IBAD deposition of thin films should provide high current density at relatively low energies in a range that, depending on the source* can span from some tenths to hundreds and sometimes thousands of eV (Ensinger, 1992). A homogeneous lateral beam profile is another desirable characteristic of these sources. Here, we will comment on the design of some typical broad beam sources used for production of low energy ions, as well as some of the more recent advances in the development of this type of device.
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
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2.4.1. Kaufmann type ion sources This type of ion source was developed by Kaufmann and is now widely used for thin film deposition (Kaufmann et al., 1989). The scheme of operation of this type of broad beam ion guns is shown in Figure 2.18. Basic elements of this source are a cathode and an anode located in a discharge chamber, a solenoid or magnetic field generator, a two-grid system separating the source from the deposition chamber and a neutraliser, usually consisting of an incandescent filament. The working gas, typically Ar or another inert gas, is introduced into the discharge chamber. Then, energetic electrons, emitted from the cathode and accelerated to the anode, strike the atoms or molecules within the discharge chamber. As a result of the electron-atom collisions, a certain number of gas atoms or molecules become ionised. While some of these ions may recombine with the electrons, mainly at the walls of the discharge chamber, other ions may pass through the holes of the first grid (i.e., screen grid) and become accelerated by the second grid (i.e., accelerator grid). The ion beam is formed by the sum of the individual beamlets produced by each hole of the aligned grids. Total beam diameter approaches the diameter of the grid system, although broadening of the beam may occur if it is not properly neutralised. To increase the ionisation probability of the gas in the discharge chamber, a magnetic field is applied between cathode and anode. The function of this field is to confine the high energetic electrons to the discharge chamber, thus limiting their recombination probability at the chamber walls. In this way multiple collisions can occur between the electrons and the gas molecules and high ion density plasma can be confined in the chamber, even for relatively low operation pressures. Owing to the high current density of the ion beam supplied by this type of source, it is necessary to neutralise the beam by injecting electrons into the beam volume. This is the function of the neutraliser, typically an incandescent filament that produces electrons to compensate the positive charge of the beam. Equal arrival rates of electrons and ions at the target surface are achieved in this way. These avoid the generation of undesired charging voltages at the target that might produce the repulsion of the incoming ions and alter their kinetic energy at the sample position. In the case of insulator materials the necessity of neutralising the beam is even more imperious than on metal since undesirable arching may be the only possibility of removing the excess of charge at the sample surface.
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L
——*
Neutralize!*
Figure 2.18. Scheme of a Kaufinann-type ion source.
Within the source, the plasma potential, and therefore the potential of its ion species, is very close to that of the anode. The ions extracted through the negatively biased grid system acquire a total kinetic energy that is equivalent to the sum of the (positive Vb) anode and (negative, Va) accelerator grid potentials. The negative potential of the acceleration grid not only contributes to accelerate the positive ions but also to repel the electrons produced by the neutraliser. Main operational parameters of this broad beam ion source are the discharge chamber pressure (or conversely the gas flow rate), the cathode emission current, the anode and acceleration grid voltages and the neutraliser current. A critical point for a proper alignment of the beam is that the holes of both the screen and accelerator grids are well aligned. Depending on these parameters, the ion current at the sample position can be modified. The distance between the source and the target also influences the actual current measured at the sample position. It is observed that the ion current decreases with increasing distance from the ion source.
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
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This is due to charge exchange between the ions and the neutral molecules in the chamber. The mean free path for charge exchange depends on pressure and path length. A typical mean free path value for a beam energy of 400 eV is about 12 mbar-cm This means that at a pressure of 10"3 mbar the mean free path is 12 cm, while for 10"4 mbar is 1.2 m. This type of ion sources is very robust and can deliver mono-energetic ion beams with a high current density in a relatively wide range of energies. They are an ideal choice for the implementation of reliable processes based on the use of Ar or other inert gas ions. However, when it is necessary to handle ions of reactive gases, this source cannot be used because the risk of burning their filaments. Other alternatives to cope with these situations should then be considered (cf., section 2.4.3). Kaufmann-type ion sources are constructed for delivering ion beams of different diameters, from some values as low as 3 cm to several tenths of cm of diameter. Typical ion densities supplied by these sources may reach values up to several mA per cm"2, depending on the acceleration voltage, usually comprised of between some hundred up to more than one thousand Volts. The low pressure of the gas required for operation, in the order of 10"3-10"4 mbar, makes these sources compatible with electron beam evaporator systems (cf., section 2.1.1) or similar devices. They constitute an optimal choice for many IBAD applications.
2.4.2. End-Hall ion sources The end-Hall ion sources are much simpler than those of the Kaufmann's type discussed in the previous section. These sources do not have any grid assembly and are very robust and reliable (Kaufmann et al., 1987). An operational scheme of these sources is shown in Figure 2.19. It consists of a filament acting as a cathode which is supplied with an alternating current, an anode at positive potential and a magnetic field produced by coils or, more typically, by a permanent magnet. When the gas is introduced into the source it becomes partially ionised by interaction with the highly energetic electrons provided by the cathode. The mixture of electrons and ions in the discharge region forms a plasma with an inhomogeneous spatial distribution within the chamber. Since the density of neutral gas molecules sharply decreases the anode downstream, most collisions with the electrons occur in its vicinity. Owing to the application of the magnetic field, plasma conductivity is higher in parallel
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than in perpendicular directions to the field lines. As a consequence, a large potential variation is found in the axial direction while the potential change is much smaller in the radial direction. Under the action of such an electrical field, the ion species are accelerated both towards the central axis of the source and towards the cathode. This latter acceleration component causes the ion species to leave the source downstream with a given kinetic energy. Meanwhile, due to the action of the radial acceleration component the ions can also cross the axis. If that happens, they can eventually be back reflected by the positive potential on the opposite side of the beam and cross the axis several times before leaving the source. This means that the beam is not collimated and diverges outside the source. This characteristic is critical for the control of the ion density at the substrate position, since it will depend on the sample position in respect to the ion source (i.e., the beam profile is not homogeneous and does not have a constant ion current). Neutralisation of the beam does not require an external filament and usually occurs through the excess of electron emission produced by the cathode. Important operational parameters of this source are pressure (or conversely the gas flow), the cathode current, the anode potential and the magnet current (when a non-permanent magnet is used). The ions generated by a grid-less source have a considerable energy spread. Significant parameters for the characterisation of the operation of a source of this type are the mean energy, the medium deviation of the ion energy (i.e., the dispersion in energy values within the beam) and the spatial dispersion of the ion current. The spatial distribution of ion energies in respect to the axial position can be approximated according to the expression: j=j0cosn6
(2.2)
where j0 is the ion current density on axis, 6 is the angle from the axis and n is a parameter that for the most common experimental set-ups ranges between 1 and 5. The end-Hall sources generate low-energy (only up to some hundreds eV), high current beams of ions. Operational parameters are similar to those of Kaufman type sources (i.e., working pressures about 10"3-10"4 mbar and ion densities values around one mA cm"2). The beam profiles, although depicting an inhomogeneous shape, are well suited for the treatment of broad area surfaces. A clear advantage is that they do not incorporate any grid assembly, thus avoiding maintenance duties related to the substitution of the grids and any possible contamination with the grid
ION ASISTBD METHODS OF PREPARATION OF THIN FILMS
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material. Owing to its low price, reliability and easy operation, these sources can be recommended for industrial applications.
Figure 2.19. Scheme of an End-Hall ion source.
2.4.3. Filament-less ion sources In both the Kaufmann and end-Hall ion sources, the plasma is generated by an electron beam produced by an incandescent filament acting as a cathode. In the Kaufmann source another incandescent filament placed outside the plasma chamber acts as neutralises of the beam. Under these conditions, only noble gases, and eYentaally nitrogen, can be used to maintain a long-term operation of the source.
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When trying to use reactive gases, such as oxygen, nitrogen or halogen containing molecules, there may be severe problems due to contamination of the sample by volatile compounds formed by oxidation and corrosion of the filaments. Moreover, a short operational time due to the burning off the filaments is another serious drawback if these sources are fed with reactive gases. Several alternatives have been developed to avoid the use of filaments to generate electrons in the plasma discharge chamber of the source (Ensinger, 1992). Among the different possibilities we can first mention the hollow-cathode ion source. In this case, the plasma discharge is sustained within the source by the electrons produced by a cold-cathode system that are extracted towards the anode by a high voltage difference. Generally, a transverse magnetic field is applied to the chamber to increase the path length of the electrons and therefore their ionisation efficiency. A two-grid system assembly similar to that incorporated in the Kaufmann-type ion sources are also implemented to extract a monoenergetic ion beam from the plasma chamber. Another means of avoiding the use of hot filaments is by generating the plasma by a radio frequency (RF) or a microwave discharge. Figure 2.20. shows a cross section of a broad-beam griddled RF ion source, operating in the MHz regime. It consists of a discharge chamber formed by a cathode and an anode, a coil to generate a magnetic field to sustain the plasma and a two-grid system for ion beam extraction. In alternative designs, a coil applies the RF field and the plasma is sustained in a quartz chamber reactor. Microwave ion sources, operating in the GHz regime, have also been applied to generate the discharge. In this case, the incorporation of suitable magnets permits working under ECR conditions with a much higher efficiency, even at relatively low gas pressures. Operation conditions of filament-less ion sources are similar to those of the Kaufmann and End-Hall ion sources. They can use either inert or reactive gases. However, the need for either RF or microwave generators makes these sources relatively more expensive than other sources only requiring simple DC or AC power supplies.
ION ASISTED METHODS OF PREPARATION OF THIN FILMS
Coil
for
87
Generator
£*rkfs
Figure 2.20. Scheme of a plasma based ion source.
References Ahmed, N.A.G., Ion Plating Technology, Developments and Applications, John Wiley & Sons, Chichester 1987. Anders, A., Surf. Coat TechnoL 93 (1997) 158. Brown, I.G., Anders, A., Dickinson, M.R., McGrill, R.A., Monteiro, O.R., Surf. Coat TechnoL 112 (1999) 271. Le Coeur, F., Pelletier, J., Amal.,Y., Lacoste, A., Surf. Coat TecnoL 125 (2000) 71. Cho, M.-H. et al, J. AppL Phys. 85 (1999) 2909. Conrad, I.E., Eadtke, J.L., Dodd, R.A., Worzola, F.J., Tran, N.C., J. AppL Phys. 62 (1987)4591.
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Ensinger, W., Rev. Sci. Instrum. 63 (1992) 5217. Espinos, J.P. et al., Adv. Mater., Chem. Vap. Deposition, 3 (1997) 219. Freeman, J.H., Chivers, D.J., Gard, G.A., Temple, W., Ion Beam Studies, Pt IV: The production of Heavy Ion Beams, Chem. Div., AERE Harwell, Oxfordshire 1977. Holber, W., Handbook of Ion Beam Processing Technology. Principles, Deposition, Film Modification and Synthesis; eds. Cuomo, J.J., Rossnagel, S.M., Kaufman, H.R., p. 21. Park Ridge, NJ: Noyes publications, 1989. Hubler, G.K, Van Vechten, D., Donovan, E.P., Kant, R.A., Mater. Res. Soc. Symp. Proc. 128 (1989) 55. Itoh, T., (ed.) Ion Beam Assisted Film Growth, Elsevier, Amsterdam 1989. Kaufmann, H.R., Robinson, R.S., Seddon, R.I., J. Vac. Sci. Technol. A 5 (1987) 2081. Kaufmann, H.R., Hughes, W.E., Robinson, R.S., Thompson, G.R., Nucl. Instr. Meth. in Phys. Res. B 37/38 (1989) 98. Khan, H.R., Frey, H., Surf. Coat. Technol. 116/119 (1999) 772. Kelly, P.J., Arnell, R.D., Vacuum 56 (2000) 159. Kelly, R., Miotello, A., Braren, B., Gupta, A., Casey, K., Nucl. Instrum. Meth. in Phys. Res. B 65 (1992) 187. Konuma, M., Film Deposition by Plasma Techniques, Springer Verlag, Berlin 1992. Lecours, A., Izquierdo, R., Tabbal, M., Meunier, M., Yelon, A., J. Vac. Sci. Technol. B 11 (1993) 51. Maiken, E.B., Taborek, P., J. Appl. Phys. 78 (1995) 541. Martin, P.J., Netterfield, R.P., Bendavid, A., Kinder, T.J., Surf. Coat. Technol. 54/55(1992)136. Marton, D., Film Deposition from Low-Energy Ion Beams, in Low Energy IonSurface Interactions, J.W. Rabalais (ed.), Wiley, Chichester 1994, p. 481. Mattox, D.M., Surf. Coat. Technol. 133/134 (2000) 517. Moseler, M., Rattunde, O., Nordiek, J., Haberland, H., Nucl. Intr. Meth. Phys. Res. B 164/165 (2000) 522. Overwijk, M.H.F, van den Heuvel, F.C., J. Appl. Phys. 74 (1993) 1762. Pulker, H.K., Surf. Coat. Technol. 112 (1999) 250. Riviere, J.P., Nucl. Instrum. Meth. Phys. Res. B 68 (1992) 361.
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Rossnagel, S.M., Methods and Techniques of Ion Beam Processes, in Handbook of Ion Beam Processing Technology, Cuomo, J.J., Rossnagel, S.M., Kaufmann, H.R. (eds.) Noyes Publ., Mill Road, N.J. 1989, p. 362. Rossnagel, S.M., Hopwood, J., Appl. Phys. Lett. 63 (1993) 3285. Sanders, D.M., Anders, A., Surf. Coat. Technol. 133/134 (2000) 78. Shedd, M., Lezec, H., Dubner, A.D., Melngailis, J., Appl. Phys. Lett. 49 (1986) 1584. Takagi, T., Yamada, I., Sasaki, A., /. Vac. Sci. Technol. 12 (1975) 1128. Voevodin, A.A., Donley, M.S., Surf. Coating Technol. 82 (1996) 199. Yamada, I., Matsuo, J., Toyoda, N., Kirkpatrick, A., Mat. Sci. Engin. R 34 (2001) 231. Zhurin, V.V., Kaufman, H.R., Kahn, J.R., Hylton, T.L., J. Vac. Sci. Technol. A 18 (2000) 37.
CHAPTER 3 EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
3.1. Ion beam effects during film growth The interest in assisting the growth of thin films with a beam of accelerated ions stems from the new properties that the ion bombardment confers to the films. In this chapter we will review the most important effects that the ion bombardment produces in IAD thin films. Most of these effects are ballistic; i.e., they are produced by the kinetic energy given up through collisions between the impinging ions and the target atoms in the most external layers of the growing thin film. To have an idea of the magnitude of the possible effects resulting from the interaction of an accelerated ion with a solid target it is interesting to remember that the energy of 1 eV per ion is equivalent to 23.08 Kcal mol 1 , when referred to a mol of ions. Binding energies associated with chemical bonds between two atoms range roughly between 100 and 300 Kcal/mol (i.e., 4.33 and 12.99 eV per single bond). Typical kinetic energies of ions for many IAD methods are in the order of 100-200 eV and above. This means that the accelerated ions have enough energy to induce bond breaking and relocation processes within the atom lattice of the growing film. A second type of effects is of a chemical nature. This effect typically appears by bombardment with reactive ion species (i.e., N2+, 0 2 + , etc.) and usually leads to the formation of oxides, nitrides or other compounds. The process parameters discussed in the previous chapter (i.e., I/A ratio, ion kinetic momentum) (cf., section 1.8), are critical for the effective control of the final properties of the film. In general, the thin film properties can be modified by a precise adjustment of these parameters. In this chapter, we will discuss some fundamental aspects of the changes induced in the composition, structure and microstructure of the films because of their growth under the bombardment with accelerated ions. For convenience, the presentation of these effects will be approached from two different perspectives. Firstly, changes incurred at a microscopical or atomistic scale. Changes in the nucleation of particles during the initial stages of deposition, roughening/smoothing of thin films surfaces and interfaces, epitaxy, mixing at 90
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91
interfaces, amorphisation or crystallisation processes, texturing of the thin films, etc., are some of the effects that, from this perspective, will be considered in this chapter. In addition, the influence of ion bombardment in modifying the intrinsic stress in thin films and improving their adhesion to the substrates will also be considered here. A second perspective stems from the analysis of macroscopic or extensive properties of the films that become modified as an effect of the ion bombardment and that define their applications and industrial use. Within this perspective, we will consider the modification of properties such as hardness, tribological properties, optical properties, electrical and magnetic properties, etc. We will deal with these potential applications in Chapter 4. Of course, both modification of the structure or microstructure and changes in extensive properties of the thin films originate from atomistic modifications in bonding and growing processes and, therefore, this distinction is somewhat artificial. In fact, the analysis at a microscopic scale of basic atom/ion interactions can be used to explain the observable behaviour of the thin films. Conversely, the determination of macroscopic properties offers a way to contrast models that describe the ultimate structure and microstructure of the films. Moreover, this distinction of properties and effects relies on differences in the experimental methodologies and techniques used to study these different aspects. Whenever possible, characterisation techniques with atomic resolution or describing the spatial distribution of atoms are used in the first instance. By contrast, in the second, measurement of observable properties or the macroscopic behaviour of the film is typically carried out. In this and the following chapter, in addition to some general concepts and models accounting for the changes induced by ion bombardment, we occasionally include a brief description of some experimental techniques used to get information about some of these ion beam effects induced in the films. However, for a more thorough explanation of the principles and applications of these methods, the reader is referred to more specialised books and reviews of thin film characterisation (e.g. Tu and Rosenberg, 1988).
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3.2. Nucleation and growth of thin films under ion bombardment The formation of a thin film implies the arrival at a substrate of atoms (i.e., adatoms when they become adsorbed onto the surface) or molecules that, by incorporation onto the growing film, give rise to its final structure and define its composition. Several processes may occur with these adatoms. One is desorption, whereby the adsorbed atoms would return to the gas phase. Association with other adatoms to form three-dimensional particles (3D) and/or two-dimensional layers (2D) may also occur. Displacements on the surface, association with surface defects or unsaturated bonding sites, etc. are other processes that have to be taken into account for a proper description of the thin film growth mechanism. As a result of all these processes, a very common stage of thin film growth is the formation of some nuclei particles with these adatoms. These nuclei constitute the seed from which the thin film will continue growing by aggregation of more adatoms and/or by association between several nuclei. The final microstructural characteristics of the thin film will strongly depend on the type, size and concentration of these initial nuclei formed on the substrate. In this section we will review the effects that ion bombardment may induce in this nucleation process and how they may alter the growing mechanism of the film.
3.2.1. Nucleation and growth of physical vapour deposited (PVD) thin films It is well established that in PVD thin films, the nucleation processes at the initial stages of the deposition are critical for the control of the growing mechanism and posterior evolution of the film. Three models typically describe the early stages of deposition of thin films prepared by supplying the material from an evaporation source (Campbell, 1997). According to the Frank-van der Merwe mechanism the deposited thin films grow according to a layer by layer process (i.e., two dimensional, 2D, growth). In this case, a first monolayer of the deposited material is formed before the second layer starts to grow. The process continues in the same way for subsequent monolayers. By contrast, in other systems the deposited material tends to form three-dimensional islands (i.e., 3D growth) from the beginning of the deposition process. This deposition mechanism is known as Volmer-Weber. Another mechanism of deposition is the so-called Stranski-Krastanov. In this case, the formation of a first monolayer of the deposited material is followed by the formation of three-dimensional islands (2D+3D growth). In principle, the occurrence of one or another mechanism depends on the surface and interface
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
93
energies of the substrate and deposited materials. However, experimental conditions such as residual pressure in the evaporation chamber or evaporation rate (i.e., kinetics of the process) may have an influence on the mechanism of deposition. In general, the most usual growth mechanism for metals deposited on insulating (e.g., oxides) or semiconducting substrates is Volmer-Weber (i.e., 3D). The size and number of the initially formed nuclei are critical factors for the control of the final microstructure of the films. Usually, a typical columnar growth mechanism follows the initial deposition steps, and the size and distribution of columns correlates with the size and surface density of the nuclei formed at the initial stages of deposition. In PVD methods, only the temperature or pressure in the deposition chamber can be effectively modified. Since a clear dependence exists between the film microstructure and the temperature of the substrate during the growing process, relatively good control of the thin film microstructure is possible in this method simply by controlling this parameter. Usually, at higher temperatures, higher atom mobility induces the formation of more compact thin films, while the opposite is true at low temperatures (see section 3.5.1 for a more detailed discussion of the mechanism of columnar growth).
3.2.2. Effects of ion bombardment on nucleation Ion assistance during thin film formation strongly affects the number, density and shape of the nuclei formed at the initial stages of deposition. Experimentally, the most commonly encountered situation is an increase in the number of nuclei particles at the beginning of the deposition process. However, the reverse, i.e., a decrease in the number of nuclei particles, has also been observed. In this case, the initial particles are larger. Nucleation is affected by bombardment, firstly because the substrate itself is modified by the ion bombardment and secondly because the island's growth mechanism undergoes different modifications under ion bombardment. We will discuss these two aspects separately. The evolution with time of the nuclei density (n) is a complex function of experimental parameters such as the energy of the ions (E), the current density (I) and the temperature of the substrate (Ts), i.e., dn/dt = J (E, I, Ts)
(3.1)
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Low ENERGY ION ASSISTED FILM GROWTH
Either with or without ion bombardment, higher temperatures of the substrate lead to an increase in the island size and a decrease in the nuclei density. This is due to an increase of the adatom mobility with temperature. The influence of the parameters associated with the ion bombardment is more difficult to predict. Ion bombardment of the substrate produces a significant amount of surface defects. These defects usually act as centres for nucleation of particles. Therefore, an increase in the nuclei density and a decrease in the size of particles should be expected at the initial stages of deposition. Since the number of induced defects increases with the ion current, an interdependence between the magnitude of this experimental parameter and nuclei density is very often found. However, during the initial stages of deposition other ion beam effects may induce a decrease in the nuclei density and an increase of particle size. Thus, it is recognised that sputtering of the deposited material may lead to the complete removal of the smallest islands. Additionally, the energy supplied by the ion bombardment may favour the diffusion of the adatoms and the dissociation of bigger clusters. Both effects tend to favour the coalescence of the deposited material into a smaller number of bigger islands. In real experiments both tendencies, leading to either an increase or a decrease in the nuclei density and to the associated result of either bigger or smaller nuclei, are acting simultaneously. Sometimes a different final behaviour can be observed by simply making slight changes in the experimental parameters. This is clearly illustrated by the example in Figure 3.1 showing a TEM analysis of the deposition of gold particles on NaCl with and without Ne+ ion assistance (Arnault et al., 1993). Through this experiment it was found that for a gold surface concentration N Au ^ 5xl0 15 cm"2 both cluster density and surface coverage are lowered for the ion-assisted sample compared with the non-assisted one. Therefore, under these conditions, cluster dissolution and/or enhanced adatom diffusivities are predominant and produce a decrease in the number of nuclei on the surface. By contrast, for NAu> 5xl0 15 cm"2, the opposite tendency is observed and a high concentration of very small clusters (covering a surface smaller than 20 nm ) is obtained. Since the ion doses increase continuously with the deposition time, an increase in the number of nucleation sites is then the predominant factor. The different nucleation and particle growth mechanism induced by ion bombardment can have a major influence on the percolation degree of the particles and, consequently, in some thin film properties, such as electrical conductivity. This effect, recognised early by Pranevicius (1979), shows that the surface conductivity
95
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
, 3
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Low ENERGY ION ASSISTED FILM GROWTH
of a growing Al film had a very long induction period before current could flow through the film. This period drastically decreased if the film growth was assisted by ion bombardment. This experiment showed that ion bombardment favours contact between a higher number of smaller particles formed under ion beam bombardment conditions.
3.2.3. Monitoring the surface defects and nucleation process induced by ion bombardment The involvement of surface defects such as nucleation centres of particles has been experimentally determined by STM. In experiments carried out on highly oriented pyrolitic graphite (HOPG), where a carbon or a nickel layer is formed by bombardment with a beam of C+ or Ni+ ions, it was shown that for ion energies above a threshold value of 38.7 eV several types of surface defects were produced (e.g. single and multiple vacancies or single interstitial between two basal planes of graphite) (Durand et al., 1998). Furthermore, it was observed that small carbon or nickel clusters form in close association with such defects. The size, shape and height of these initial nuclei could be monitored using this technique, as could their posterior evolution when increasing the ion doses. For a similar experiment, Figure 3.2 evidences the influence of surface defects in the formation of nuclei. It shows two STM images of Ni particles deposited on HOPG under different experimental conditions. The two images have
(a) Ni-13A1L1 (Ni on virgin HOPG)
(b) M-20F2L1 (Ni onAr* induced defects)
Figure 3.2. STM (100x100 nm) images of Ni films deposited by: (a) evaporation on a virgin HOPG substrate; (b) evaporation on a 100 eV Ar+ irradiated substrate. Reproduced from Durand et al. (2000) with permission.
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been taken for equal amounts of Ni deposited by evaporation on a virgin substrate (a) or on a substrate that has been previously subjected to 100 eV Ar+ bombardment (b). It is clear from the comparison of the two images that on the bombarded substrate the number of nuclei particles increases while their size decreases with respect to the situation of the virgin substrate. This difference must be attributed to the development of nucleation centres around the surface defects generated by the initial ion bombardment treatment.
3.2.4. Description of nucleation and growth processes by analysis ofSTM/AFM images A deeper insight into the growth mechanism of a thin film in its early stages of formation can be obtained by applying the scaling theory to the AFM/STM images of a growing layer (Family et al., 1985). An interesting parameter that can be derived from the STM images is the so-called Dynamic Scaling Function of Roughness (DSFR), a. The usual method to calculate the DSFR divides each STM/AFM image (e.g., formed by 256 x 256 pixels) in smaller images of length L (e.g., of 128 x 128, 64 x 64, ..., 2 x 2 pixels). The specific roughnessCT(L,t), that is the roughness for a specific length scale L at a time t (or, what is equivalent, for a given amount of deposited material), is then obtained by calculating the root mean square (RMS) roughness inside each L x L image and averaging over the ensemble of images of the same size. In this way each image is reduced to one-dimensional function. According to this theory, a given DSFR is characteristic of each type of growth mechanism. For many growth models, the DSFR log-log plots with respect to the specific length (i.e., log a vs. log L) give two distinct regions separated by a crossover length L0. The slope at each point of this log o/log L function, %, is called the roughness scaling exponent and its value for L0 and whose specific value is a measure of the efficiency of the diffusion processes. By contrast, for the region L > Lo, the adsorption/desorption terms are monitored, the slope is % - 0 and the corresponding a value is equivalent to the average roughness of the film. These two regions, where %*0 and %=0, are clearly discerned in the curves presented as an example in Figure 3.3. This figure shows the log-log plots of
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Low ENERGY ION ASSISTED FILM GROWTH
the DSFR function values deduced from STM images similar to those in Figure 3.2. In this case, the experiment consisted of the formation of a carbon layer on HOPG by direct bombardment (IBD, cf., section 2.2.5) with lOOeV C+ ions (Durand et al. 1998). Each plot in the figure, deduced from the analysis of a given STM image, corresponds to sample situations of increasing ion densities. From a first assessment of the plots it is clear that for the region L
Dynamic Scaling Function
1
10 100 Specific Length L (nm)
1000
Figure 3.3. Evolution of the Dynamic Scaling Function of Roughness as a function of ion density for the early stages of growth of a C film deposited on HOPG by bombardment with 100 eV C ions. Reproduced from Durand et al., (1998) with permission.
mechanism of surface growth. It reveals that with high ion densities relatively less adsorption defects are formed, and that adatoms are more likely to assemble at the point of an already existing defect-particle association than remaining in their own position. On the other hand, for L>Lo the thin film roughness reaches its maximum value for densities around 15-20 ions/nm2 to slightly decrease for higher ion dose. This tendency has been interpreted by assuming that for very high ion densities the growth of islands is the prevailing factor controlling the average thin film roughness. Under these conditions, the maximum roughness corresponds to the maximum height of the islands. Above this peak value there is a limitation to the height of islands before they collapse or meet with a neighbouring island and a steady state roughness is reached whatever the ion dose.
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99
3.3. Topography and surface and interface roughness One of the advantages of assisting the growth of thin films by ion bombardment is that their microstructure can be modified to improve certain properties that are beneficial for the intended applications of the films. In this section we will review some microstructural and morphological changes that can be induced in IAD films. The changes refer to modifications in crystal size and in surface and interface roughness. In section 3.5 we will refer specifically to the densification of IAD films, perhaps one of the most interesting properties of this type of thin films when compared with others prepared by evaporation methods. Proper control of the thin film properties may be critical for their mechanical, optical, electrical or magnetic properties, as it is indispensable in many cases to use IAD processes to get films with the adequate characteristics. As examples that illustrate the great importance of preparing flat thin film surfaces or interfaces, let us mention how critical it is to produce thin films with a low roughness and small grain sizes to decrease the light scattering in optical surfaces or for an efficient control of exchange coupling phenomena in magnetic multilayer systems (cf., Chapter 4).
3.3.1. Grain size It was recognised during the early stages of the development of the IAD techniques that a decrease in crystal size usually occurs when the growth of the thin films is assisted by bombardment with low energy ions (Smidt, 1990). Decrease in grain size is a factor that can contribute very efficiently to the densification of the films. The effect of ion bombardment in decreasing the grain size has been associated with an enhancement of the nucleation rate during the film growth. In section 3.2, the fact that one of the possible effects of ion bombardment during the initial stages of surface nucleation is to increase the number of nuclei by producing surface defects that act as nucleation centres was discussed. In some particular cases it has also been observed that ion bombardment may promote an increase in crystal sizes. This situation would indicate that other phenomena leading to an increase in adatom mobility, due to higher local temperatures or high strain energy, might favour crystal growth. In IAD thin films the crystal size is very sensitive to the beam energy and to the I/A ratio. Thus, for metal thin films a sharp decrease in crystal size for relatively low beam energies up to 60/100 eV is generally observed. The size is
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Low ENERGY ION ASSISTED FILM GROWTH
not significantly modified for higher energies, provided that the beam energy does not reach values above 10 keV. In fact, for beam energies of the order of some tenths of keV, it is very common to find higher grain sizes. High-energy ions can penetrate deeper into the film and may influence recrystallisation and grain boundary motion in buried regions, where grain size increase may then occur.
3.3.2. Surface roughness One of the expected morphological effects of a decrease in grain size during IAD growth is a decrease in the roughness of the surface of the film. In fact, this is the expected result when small grains form at the surface since, at a first approximation, the roughness of a thin film can be considered equivalent to that resulting from the coalescence of the growing particles that form the film. However, in practice, other factors besides island growth, coalescence and grain size, have to be considered to properly account for the surface roughness of thin films grown under ion bombardment. Smoothing or roughening of thin film surfaces are very complex processes that depend on several factors resulting from the interaction of the energetic ions with the growing film. While some of these additional effects may contribute to the roughening of the thin film surfaces, others contribute to their smoothing. In practice, although in most cases ion bombardment during growth leads to smoother surfaces, there are cases where roughening has been also reported. Basic processes that have to be considered to account for the final thin film roughness are the increase of adatom mobility and the occurrence of sputtering phenomena. In general, an important effect contributing to surface flattening is the increase in adatom mobility induced by the transfer of energy from the ion beam to the growing surface. Thus, low-energy ion bombardment would favour local atomic rearrangements enabling the adatoms to relax into low energy sites at step, terrace or kink positions and, in this way, contribute to surface or interface planarisation. In this respect, it is expected that for equivalent experimental conditions, surface roughness will depend on the mobility of the depositing atoms and that, therefore, higher surface roughness should be expected for the deposition of materials with high activation energies of diffusion. Molecular dynamic simulations have been used to simulate flattening of deposited metal clusters under ion bombardment (Zhou et al., 2000). The results show that the flattening degree, estimated from the number of adatoms in contact with the substrate, depends on the type and energy of
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the ions, the use of Xe instead Ar ions for producing flat surfaces being more efficient. This result has to be associated with the higher momentum transferred by the heavier Xe ions (cf., section 1.8.3). Sputtering is another effect occurring during the ion beam bombardment of a material. Sputtering may contribute to the roughening or flattening of the solid surfaces. Carter (1998) has formulated deterministic smoothing and roughening processes and developed several differential equations to describe the evolution of surface morphology of growing thin films. This analysis predicts that roughening may occur as a result of a variety of sputtering phenomena. Thus, in the absence of any other smoothing or atomic relaxation effect, the stochastic character of the sputtering should contribute indefinitely to surface roughening. However, on a microscopic scale, this tendency can be counterbalanced by different factors contributing to either smoothing or roughening and that they compensate each other. In fact, the sputtering yield is a function of the orientation of the incident ion flux with respect to the surface normal (cf., section 1.7.1). This effect tends to preferentially erode any developing surface tip or asperity and, in this way, contribute to smoothing. An opposite roughening effect stems from the fact that the sputtering yield is larger for troughs than for asperities. An additional sputtering effect leading to smoothing is the creation of recoil atoms parallel to the surface. These recoil atoms moving on the surface tend to maximise neighbour bonding between atoms by saturating low-coordinated atom sites. In this way, they may contribute to a curvature-dependent smoothing process since the concentration of coordination vacancies will be higher for rougher and more curved surfaces. It is also generally recognised that the amorphous character of the deposited material favours the smoothing because in this case a viscous flow mechanism for transport of material on the surface may have a significant importance in decreasing the surface roughness. Present knowledge makes it difficult to predict what the final effect of the ion bombardment on the roughness of a thin film will be. However, there are some empirical observations that can be used to make predictions about the evolution of surface roughness as a function of the actual values of some experimental parameters. It has already been mentioned that a common pattern of behaviour observed under many experimental conditions is an enhancement of roughening for ion bombardment at significant off-normal angles with respect to the surface normal. By contrast, smoother thin film surfaces are commonly obtained for normal ion bombardment. Other experimental parameters that can affect the smoothing are
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Low ENERGY ION ASSISTED FILM GROWTH
the ion energy and the ion to atom ratio (I/A, section 1.8.2). A decrease in surface roughness is observed for low ion beam energies and high I/A ratios and, in general, when ion flux density approaches the atomic deposition rate or film-growth rate. A problem occurring in this case is that the final thin films turn out to be relatively thinner because of sputtering erosion. Direct determination of surface roughness is possible by AFM. From the analysis of the AFM images, it is possible to estimate a RMS parameter to measure the thin film roughness. As an example, Figure 3.4 shows the dependence of this parameter for Ta2Os thin films grown by DIBS. It is apparent in this plot that ion bombardment produces a sharp decrease in surface roughness for beam energies between 50 and 150 eV. Meanwhile, at higher beam energies, the surface roughness slightly increases, probably because some sputtering processes inducing roughening are more important at these energies (Lee et al., 1997). In this experiment it is also worthy to note that surface roughness was very sensitive to the composition of the ion beam. Minimum roughness was found for a beam composition characterised by an 0 2 /Ar ratio around 0.4/0.6. Lower or higher ratios yielded higher values of the RMS parameter. Incorporation of Ar atoms within the lattice (cf., section 3.6.4), amount of momentum transferred to the growing surface (cf., section 1.8.3) or other subtle effects related to the presence of Ar in the beam must have some influence on
0.5
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Ion beam voltage of the second ion source,Vb2 (V) Figure 3.4. Surface roughness vs. energy of the assisting ions for Ta 2 0 5 thin films grown by DIBS. Reproduced from Lee et al. (1997) with permission.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
103
the evolution of surface roughness during growth and, therefore, be critical for the determination of the final planarity of the films.
3.3.3. Step and surface coverage Improvement of step and surface coverage and planarisation of surfaces are some of the beneficial effects of ion bombardment in IAD thin films. Surface and step coverage can be improved by ion assistance, even if the zones to be covered are under a shadow area with respect to the evaporation sources. Several factors can contribute to such homogenisation of the distribution of the deposited material when the film is being subjected to ion bombardment. One is the increase in surface mobility of adatoms when the film is ion bombarded. Sputtering from some places and redeposition in others of the thin film material is an additional effect contributing to the redispersion of the deposited material. Dependence of the sputtering yield with the angle of the incoming ions at a particular point of the surface can contribute to surface planarisation since the film will grow further in those zones where the sputtering yield is smaller (cf., section 1.7.1). Usually, ions are incident perpendicular to the macroscopic substrate surface, but oblique at microscopic features, steps, etc. of the surface, contributing differently in each zone to deposition and/or sputtering phenomena. Planarisation of thin films, even if the substrate presents a rough surface and perfect coverage of steps and grooves in zones out of sight of the evaporation source are useful effects for practical applications. This feature is particularly interesting in the microelectronic industry and for the more recent micromachining developments, where complex structures have to be homogeneously covered by the thin film.
3.3.4. Surface roughness of thin films grown by IBD Previous considerations apply to thin films whose growth is assisted by bombardment with inert gas or oxygen or nitrogen ions supplied by an independent ion source. In several methods of deposition of thin films, accelerated species of the thin film material impinge directly on the substrate surface (cf., section 2.2). For
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Low ENERGY ION ASSISTED FILM GROWTH
these types of deposition procedures, additional ion induced effects have been reported to contribute to surface smoothing. Carbon thin films can be prepared by direct C+ ion bombardment (cf., section 2.2.5). For this synthesis procedure it has been found that below a certain energy of the ion beam, the thin films show rough surfaces and have a graphite character. By contrast, for ion energies higher than 30 eV the thin films become very smooth and the films have a diamond-like character. The sp2 or sp3 character of the carbon bonds in the films is the criterion that is typically considered for the evaluation of the relative graphitic or diamond-like character of the film (for more details see Chapter 5). The so-called subplantation model by Lifshitz et al., (1994) (cf., section 5.12.3) accounts for these observations. Within this model, it is considered that with energies higher than 30eV, the carbon atoms may penetrate the thin film surface and become incorporated in subsurface positions. The subplantation of these atoms provokes a high internal stress and the formation of a dense diamond-like phase. Meanwhile, the growing surface of the film may retain the initial smoothness of the substrate because no layer is growing on the surface, it develops embedded under the first substrate layers instead. By contrast, for energies lower than 30eV, most carbon atoms remain on the target surface, where they tend to form nuclei that coalesce and grow as a graphite thin film. In this case, roughness is the result of the aggregation of material on the substrate surface as in a typical PVD process.
3.3.5. Interface roughness The development of sophisticated optical and electronic devices requires the use of thin film or multilayer structures where a very precise control of the thin film/substrate or layer/layer roughness is required. Even for multilayer structures prepared for mechanical applications, the control of the interface characteristics is very important. A particular case that illustrates the importance of such interface control is that of epitaxial thin films where, at the interface, the atoms of the film have to be in register with those of the substrate. An effective control of the interface smoothing during IBAD can be achieved by an adequate choice of the deposition parameters. As has been discussed previously, when dealing with the surface roughness of thin films, ion energy, I/A ratio and type of projectile have a definitive influence on the interface quality of these systems. Molecular Dynamics has been used to simulate the evolution of interface roughness with these experimental parameters (Zhou et al.,
EFffiCTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
105
2000). Figure 3.5 shows representative atomic structures of an Ni/Cu/Ni multilayer grown by evaporation of the two metals and simultaneous bombardment with Xe+ ions of increasing energy. The simulations clearly show that the interfaces become smoother when the growth of the film is ion beam assisted. An increase in adatom mobility is the main reason for an enhancement in surface and interface smoothing. The simulations in Figure 3.5 also reveal that for higher ion energies some interface mixing may be induced, an effect that can be deleterious for certain applications. In practice smooth and well defined interfaces in IAD thin films or multilayers are produced by using relatively low ion energies, typically below 50 eV, and relatively
(a) Notonbombardment
-*• x t « i f
(b)E».0.boV
(C) EXe ^3MeV
Figure 3.5. Representative atomic structures calculated by molecular dynamics simulations of Ni/Cu/Ni multilayers as a function of assisting Xe ion energy at an ion/metal atom ratio of 2, a metal atom energy of 0.1 eV, a substrate temperature of 300 K and a normal incident angle; a) without Xe ion assistance; b) ion energy of 0.5 eV; c) ion energy of 3.0 eV. Reproduced from Zhou et al. (2000) with permission.
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low I/A ratios. These experimental conditions aim to increase the adatom mobility, while keeping to a minimum undesired effects such as interface mixing (subplantation of C into graphite may occur already for energies above 38 eV) (cf. section 3.3.3) or sputtering induced roughness that are more important for high energy ion beams. Another beneficial effect of ion bombardment consists of the suppression of the replication by the first layers of the large-scale lateral roughness existing in some substrate surfaces. This effect has been observed on Co/Cu multilayers prepared by low energy ion-assisted deposition where the successive layers were progressively more dense and smoother than the substrate, while they did not present any replication of the substrate roughness (Telling et al., 1998).
3.3.6. Monitoring the interface roughness by X-ray reflectometry Direct observation by TEM is a well known method of proving the quality of an interface between a thin film and the substrate or between the different layers of a multilayer structure. Many examples can be found in literature about the potentialities of this technique for observation of interface quality (De Hosson et al., 2001). However, for practical applications it is sometimes necessary to determine the average roughness over large interface regions, rather than obtaining information restricted to the field of observation of TEM. A suitable technique for the determination of the average roughness at surfaces and interfaces is X-ray reflectometry (Nevot et al., 1980). Besides surface and interface roughness parameters, this technique provides information about the electronic density (which is directly correlated with the atom density of the material and therefore gives information about the compactness of the film) and thickness of the examined layers. X-ray reflectometry spectra are collected by irradiating the thin film surface at very grazing angles below and above the total reflection angle in the X-ray region of the investigated material. The reflected intensity of the X-rays measured as a function of the incident angle gives a typical interference spectrum that can be simulated with a proper model of the thin film structure. Figure 3.6 shows as an example the experimental and simulated reflectivity curves for a TiN thin film prepared by DIBS (Alvisi et al., 1997). The interference pattern (kiessig fringes) results from the interference of the X-ray waves reflected at the air-film and film-substrate interfaces. In this way information about the interface thickness and roughness can be obtained. This technique can be used for characterisation of thin
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107
films with a thickness able to give rise to interference patterns. For the majority of materials, values around 100 nm can be considered as the highest limit of the thickness of the thin films that can successfully be investigated with this technique. The experimental curve in Figure 3.6 can be well simulated by assuming a thin film structure as that schematically shown in the right part of the figure. The proper simulation of the interference patterns requires the assumption of three values of interface roughness and the corresponding density parameters for the film and interface zones. The best fitting parameters are taken as the actual values of the corresponding magnitudes for the system under investigation. In the present case, besides the central layer of TiN, the assumed layer structure consists of an outer layer of Ti0 2 and a substrate-film interlayer of SiC>2. This structure is realistic since the samples are exposed to air and their surface is likely to be oxidised to Ti0 2 . Meanwhile, the substrate is likely to develop a thin Si0 2 interlayer formed during preparation of the film by oxidation with the residual gases of the deposition chamber and eventual mixing with the material of the thin film.
2nm
^
0"i=1.5nm .O"2=0.7nm
35nm
1.5nm
^—0"3=1.3nm
Figure 3.6. (Left) Calculated (upper) and experimental (lower) reflectivity curves for a 35.0 nm thick TiN film grown on a Si(100) substrate. (Right) The schematic structure of the thin film refers the layer structure and the parameters used for the X-ray reflectivity calculations. Reproduced from Alvisi et al. (1997) with permission.
Besides the determination of the interface roughness, whose values are reported in the scheme of Figure 3.6, it is also interesting that the reflectometry analysis of the 35 nm thick TiN layer provides information about its electron density and, conversely, about its atom density. The electron density obtained in this case was very high (i.e., 1.53-103 nm"2), in agreement with the very compact and dense film structure obtained by using IAD procedures.
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3.3.7. Epitaxial growth of thin films IAD deposition methods are also used for growing epitaxial thin films or multilayer structures. The use of ion beams is becoming an advanced molecular beam epitaxy (MBE) method where additional energy for the layer growth is provided by accelerating all or a part of the incoming species arriving to the target. The beneficial effects of ionising and accelerating some of the incoming species has been clearly shown by the growth of epitaxial Si layers on a (100) oriented Si substrate by MBE of partially ionised and accelerated Si+ species (Wagner et al., 2001). In this homoepitaxial deposition of silicon on silicon, mixing effects are obviously neglected. The advantages of the ion beam assistance in favouring the epitaxial growth are not restricted to homoepitaxial systems. Thus, heteroepitaxial growth of GaN thin films on c-plane sapphire (i.e., A1203) single crystal substrates has been carried out by nitrogen ion bombardment during Ga evaporation (Gerlach et al., 2000). The crystalline quality of the epitaxial thin film was very much dependent on ion beam characteristics, the highest crystallinity being obtained for beam energies below 50 eV. Again, it is interesting to remember here that low energy values of the ion beams (i.e., below -30 eV) are required to avoid mixing or sub-plantation of impinging ions beneath the first substrate layers (cf., section 5.12.3).
3.4. Interface mixing Atom displacements induced by ion-target collisions may cause a modification of the in-depth distribution of the elements. Mixing effects have as a consequence that the atom distribution does not present sharp profiles at the interface as could be expected for PVD films. This is due to the mixing of layer and substrate atoms (or atoms of two layers when dealing with multilayer structures) that after ion impacts become distributed in the substrate and layer, respectively. While interface mixing has to be reduced to a minimum for the growth of epitaxial or sharp multilayered structures, for other applications interface mixing is advantageous and is promoted to obtain thin films with good adhesion to the substrate. In fact, mixing enhances the adhesion between thin film and substrate, so that delamination can be avoided if an effective atom redistribution occurs at the interface. Basic macroscopic aspects of adhesion will be treated in section 3.11. Here, we would like to stress that in many cases, such as in metal/ceramic, metal/polymers or oxide/polymers systems, atom mixing at the interface is the best procedure to achieve good adhesion. In thin films
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
109
prepared by IAD methods it is possible to get such a mixing already during preparation of the films without any post-deposition treatment (i.e., thermal annealing, etc.), just by selecting the appropriate ion bombardment conditions. In this respect, a significant advantage when trying to deposit metals or oxide thin films on polymers is that deposition by IAD methods is possible at room temperature, a requisite imposed by the stability of the substrate.
3.4.1. Mixing in thick films and bulk materials induced by high energy ions Ion beam mixing in thick films and bulk materials has been widely studied for more than three decades, either experimentally or theoretically (Paine et al., 1989). For these investigations relatively high ion energies (i.e., various tens of keV) have been utilised. Applications which are currently in use and those evolved from these studies are the synthesis of new materials by mixing buried layers or multilayer structures of different compounds. Temperature, ion energy and fluency and mass of ions are effective factors for the control of the extent and efficiency of mixing. In principle a detailed account of this subject is not the subject of this book, mainly dedicated to IAD thin films. However, a brief outline will be included here because some of the basic concepts may be of some interest when using relatively highenergy ions for the synthesis of the films. Mixing effects have been primarily interpreted in terms of ballistic processes: i.e., by considering the efficiency of energy transfer by collisions between the impinging ion and the target atoms, and the consequent atom displacements that derive from such events (nuclear and electronic stopping efficiency) (cf., section 1.3.1). However, it was quickly recognised that the extent of mixing was also a function of the type of materials brought into contact. This material specificity suggested that mixing is a chemical driven process too and that thermodynamic and kinetic factors controlling bond formation should also be taken into account in explaining the formation of new phases during ion mixing. The thermal spike model has been successfully applied to describe chemically driven mixing effects (Mayer et al., 1981). In this model, it is assumed that the local energy deposited in a collision cascade is very high and that, therefore, the local temperature in the volume surrounding the ion tracks may reach very high values. In this way, ion beam interactions can produce similar chemical phases to thermal treatments. By contrast, it is not so clear whether significant thermal transport of mass can occur in the short quenching times (~1 xlO"11 s) following the primary
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Low ENERGY ION ASSISTED FILM GROWTH
thermal spikes produced by the ion impacts. Thus, it is generally admitted that the mixing processes mainly occur before the energy accumulated in the tracks is thermalised and are a result of the ballistic interactions between the accelerated ions and the target atoms leading to atom displacements out from their original positions. In section 1.8.3, a brief comment was made regarding the influence on mixing of the magnitude of the ion momentum transfer during ion bombardment.
3.4.2. Interface mixing in IAD thin films Interface mixing is commonly observed in IAD thin films, even if the value of the ion energies is much smaller than in the aforementioned experiments with buried layers. The ion-solid collision theories are not very accurate for E<100 eV and no proper theoretical descriptions of mixing are available at present. Therefore, most mixing studies in this area rely on experimental evidence. As a rough approximation the thickness of the mixing layer W can be considered as proportional to the square root of the product of the ion dose <j) and the nuclear stopping power: W~(Sjn
(3.2)
Eqn. (3.2) is a simplification of the problem because it has been deduced under the assumption that only ballistic effects are involved in the ion-target interaction. From a practical point of view, this expression indicates that in IAD experiments, when the ion current is set fix, the thickness of the mixing layer can be considered proportional to the square root of the ion mixing time. However, it has been found experimentally that the extent of mixing also depends on other variables, such as the beam energy, the I/A ratio, the type of ion and the chemical nature of the two components brought into contact at the interface. Molecular dynamic simulations have been carried out to describe the mixing processes occurring at the interface during IAD thin film formation (Nordlund et al., 1998). An interesting result from these calculations is that not only is the chemical nature of the components brought into contact important to define the mixing efficiency, but also the crystalline or amorphous character of the materials. In fact, compounds with amorphous structures seem to be more favourable for mixing than compounds with close-packed structures. This difference has been associated with the fact that low energy atoms may migrate more easily
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111
through open and soft networks than through compact and well-crystallised structures. To be able to make comparisons between different experiments it is customary to use a mixing efficiency function that is defined as: Q = Dt/<j>'FDN
(3.3)
where D is an effective diffusion coefficient for mixing. It can be determined from experiments with marker systems consisting of a thin layer of an element or compound sandwiched in another material (Paine et al., 1989). t is the bombardment time, (j)' is the ion fluency and FON the deposited nuclear energy per ion and per unit depth. FDN values can be calculated and are conveniently tabulated (Winterbon, 1975). The mixing efficiency function enables the comparison between different experiments or between experiments and theoretical results, since it depends on experimental parameters and on the chemical nature of the elements brought into contact.
3.4.3. Monitoring interface mixing by TEM/EELS Interface mixing during IB AD preparation of thin films can be determined by direct observation and analysis of the interface by means of transmission electron microscopy (TEM). X-ray reflectometry can be also used to determine a macroscopic roughness parameter for the interface that, in turn, can be related with the degree of interface mixing. The combined use of TEM plus the electron energy loss analysis (EELS) of a line profile along an interface may provide a clear view of the mixing processes that occur at an interface during preparation of a thin film (Sohn et al., 2000). Figure 3.7 shows as an example two cross-sectional highresolution TEM images for two carbon films grown on Si (100) by C" ion beam bombardment with energies between 300 and 500 eV. The images reported in this figure show that the interface is severely damaged when using 500 eV C ions, while it is not very much disrupted with 300 eV C" ions. In the latter case the interface is smooth and a well ordered silicon network is clearly seen throughout the interfacial area. The possibility that C-Si bonding formation occurs at the interface can be investigated by following the evolution of the Si L2,3 EELS spectra along a line crossing the interface. This technique also provides quantitative information about the change of the C/Si ratio along the interface. The inset in Figure 3.7 shows a
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Low ENERGY ION ASSISTED FILM GEOWTH
series of L23 edge spectra taken from the ieterfaciai region for a film grown with 500 eV C ions. This series of spectra can be attributed to metal silicon (Si L23 first peak at 100.1 eV) and to silicon in SiC (Si L23firstpeak between 103.5 and 104.5 eV). In this way, EELS characterisation of the interface demonstrates that carbon was not only mixed with silicon at the C/Si interface as suggested by the TEM micrograph, but also that carbon has formed chemical bonds with the silicon atoms within a substeate region of 2-3 ran. The existence of this intermixed C/Si layer in the 500 eV sample was critical for a good adhesion of thefilm,since for the sample prepared with 300 eV C ions thefilmdelaminated very easily.
Figure 3.7. Cross sectional TEM micrographs of an amorphous carbon film grown on Si (100) substrate viewed along the [110] axis, a) The substrate was bombarded with a 500 eV C* beam, b) The substrate was bombarded with a 300 eV C beam. The inset shows the Silicon L-edge spectra collected at the C/Si interfacial region of the 500 eV modified sample. The interval between each spectrum is 2.1 em. Reproduced from Sohn et at. (2000) with permission.
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3.5. Densification of thin films IAD thin films usually present a high density that, in some cases, may even approach values typical of the corresponding single crystal compounds. This high density is recognised as one of the most characteristic features of thin films prepared by IAD methods.
3.5.1. Columnar growth in PVD thin films Thin films prepared by PVD present a columnar microstructure with a high concentration of voids between columns or within the columns themselves. The open space in these thin films can be a significant fraction of the total volume of the layer and may vary within wide limits depending on the temperature of the substrate, the growth rate, the residual pressure in the chamber and other experimental factors. The columnar structure of PVD films is basically the combined consequence of two factors, the shadowing effects produced by the first grains formed from the initial nuclei and the low adatom mobility under the conditions of deposition. Shadowing prevents the vaporised atoms coming from the source to reach certain zones of the substrate, while their low mobility favours their staying in their arrival positions at the surface rather than moving to occupy the void zones. Temperature and pressure during deposition may strongly affect the characteristics of the columnar growth and the void volume fraction within the film. Such a relationship between microstructure and processing conditions of the film can be viewed in a simplified structure zone diagram (Thornton, 1977). A schematic representation of such diagram as a function of the deposition temperature (Ts) referred to the melting point temperature (Tm) of the material is reported in Figure 3.8. In this diagram different zones are recognised where the crystal and columnar growth mechanisms are controlled by different processes as indicated in the figure (i.e., surface diffusion, grain growth, etc.). In the middle zone II, a typical columnar growth takes place with a significant fraction of void volume between the columns. A clear effect of increasing the deposition temperature in zone III is an increase of the crystal size because of the enhancement of the adatom mobility. The bigger size of the crystals and columns in this zone means that the overall void volume fraction between them decreases with respect to zone II. Moreover, the higher adatom mobility in this zone enables the filling of some void spaces, thus leading to an
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increase in thin film density. In the I and T zones, the small kinetic energy of the adatoms prevent their diffusion on the surface, so that they practically remain at or close to the positions where they arrive at the surface. In this zone the thin film grows by coalescence of a large number of small crystallites.
Zone I
Zone T
Shadowing
Zone II
Zone III
Surface diffussion
Bulk diffussion; grain growth
Figure 3.8. Evolution of the type of columnar growth as a function of substrate temperature for PVD thin films
3.5.2. Densification in IAD thin films Ion bombardment during thin film growth produces, as an effect, a considerable increase of adatom mobility that, together with other effects that will be discussed latter in this section, provokes a disruption in the columnar growth of the films. Densification is a first microstructural consequence of these changes in the thin film growth mechanism. Another typical effect is the decrease of crystal size, a phenomenon that contrasts with the increase in crystal size observed in PVD films when the substrate temperature increases. Such effects mean that, in general, the microstructure of IAD films falls within zone I of the microstructure diagram in Figure 3.8. In IAD thin films the crystallites become very effectively packed because the enhanced adatom mobility induced by the ion bombardment. A minimum formation of voids or pores between the small crystallites is a differential effect of the densification induced by the ion bombardment.
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The main reason which accounts for densification in IAD thin films is the removal of voids during thin film growth because of enhanced adatom mobility. A possible mechanism for the transfer of ion energy to the target atoms could be that of a thermal spike (cf., section 1.6.2). However, detailed calculations have shown that densification would not occur as resulting from this type of processes. Miiller (1986) and Netterfield et al. (1988) proposed a model based on collisional cascades to explain the enhanced densification of IAD thin films. According to this model, momentum is transferred from the incoming ions to the target atoms while, simultaneously, a series of secondary effects including sputtering, recoil and atom implantation, and diffusion of target atoms from non-stoichiometric regions in the case of oxide samples, are also induced. The contribution to density enhancement of the above-mentioned effects can be visualised by molecular dynamic simulations (Miiller, 1987). Calculations with a 2D structure of Ni atoms show that atom rearrangements produced in a collision sequence lead to a collapse of voids and to a transport of atoms from one point to another of the surface. Thus, it was shown that while recoil events of subsurface atoms may serve to fill some internal voids, the increased mobility of surface atoms results in the filling of surface holes. Both effects lead to a more densely packed microstructure. In addition, the enhanced adatom mobility would also favour the smoothing of the surface. The influence of substrate temperature and kinetic energy of the ions on the void size and thin film density, as well as the tendency of voids to align in vertical tracks has also been simulated by molecular dynamic calculations (Smith et al., 1996). Figure 3.9 shows the results of a calculation that illustrates the influence of the ion energy (expressed in terms of reduced energy e, cf., section 1.2.4) on the growth process of a thin film under ion assistance. It is apparent in this figure that void volume drastically decreases with the energy of the impinging ions. Comparison with similar calculations as a function of the deposition temperature shows a similar trend, thus suggesting that target temperature and kinetic energy of the impinging ions may similarly influence the deposition process. Another interesting feature deduced from these calculations is that the voids remaining in the film define ideal lines which are parallel to the ion beam direction and that connect with some surface imperfections. This interesting result has been explained by reproducing the formation of voids through a mechanism that involves surface holes. When new atoms and ions arrive at the surface there is a certain probability that an atom bridge forms between the walls of surface holes, without the inner part of the hole being filled with new atoms. When a bridge is
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Low ENERGY ION ASSISTED FILM GROWTH
successfully formed a new void remains embedded within the film in the same vertical location where the hole initially occurred. Successive voids can form in the same direction because the surface imperfection is not completely removed by the formation of the bridge and the process may happen again at the same position at a latter stage of the thin film formation. Of course, with a sufficient supply of energy, this process may lead to a complete filling of the hole and to the removal of any void and surface imperfection.
(a) 5,= 0.10 e
0>) ^,= 0.45e
(d) Ej,= 1.15 s
(c) E b =0.80e
(e) § , = 1.50 8
Figure 3.9. Typical microstructures for films grown at various ion energies (Eb) onto a substrate at a temperature T=0.125 e/k. Calculations correspond to a tridimensional structure of Ni atoms (e= 0.74 eV). Reproduced from Smith et al. (1996) with permission.
3.5.3. Evolution of density and crystallinity with ion energy and I/A ratio Energy and I/A ratio are critical parameters for an enhancement of density in IAD thin films. Experimentally, it is generally found that any increase in the value of any of these two parameters leads to an increase in the density of the film up to reach a
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
117
maximum value. Then, for higher values the density remains constant or decreases slightly. This tendency is schematised in Figure 3.10 for the I/A ratio. The initial increase in density resulting from ion beam bombardment has to be associated with the disappearance of macroscopic voids. Above a certain limit, where most of these voids have already disappeared, other processes start to contribute to a decrease in density. In fact, a further increase of the ion current may induce gradual increases in structural disorder and damage and/or the generation of crystallographic defects. This is illustrated in Figure 3.10 by the curve representing the crystallinity of the film as a function of the I/A ratio. The maximum of the crystallinity and density curves appear at similar I/A values. This common evolution is logical since any increase in adatom mobility should result in atom rearrangements leading to both higher packing density and better structural order. Above the I/A value of the two maximum, accumulation of defects would lead to a progressive decrease in crystallinity, a feature that, for an extreme situation, might result in the complete amorphisation of the lattice. The influence of the accumulation of lattice defects on the density is less pronounced than that of the removal of voids and, therefore, the density only decreases slightly after the maximum.
P bulk
"E 75
3
O
l/A
•
Figure 3.10. Influence of the I/A ratio on the density (p) and crystallinity of IAD thin films.
Experimental investigations looking at variations of thin film properties such as refraction index that are very much dependent on thin film density have confirmed the previous behaviour scheme (cf, section 4.5).
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Low ENERGY ION ASSISTED FILM GROWTH
In general, an increase in density is also found as the energy of the ions increases. Some models have been developed to estimate the optimum average energy that has to be supplied per deposited atom to achieve a film with the maximum density. These models take into consideration that maximum packing density can be achieved by assuring at least one displacement per condensed particle. The calculated average energies range between 20-30 eV for oxide and nitride materials to 60-80 for metals (Grigorov et al., 1988). For these energies both ion penetration and particle rearrangements would take place at the surface of nearsurface layers while the ion induced defects in these regions could be annealed very efficiently at low deposition temperatures. These predictions rely on the assumption that only the overall amount of energy and momentum transferred to the target atoms are important for the enhancement of the thin film density. However, many experimental results have pointed out that, depending on energy ranges, the actual value of the ion energy rather than the average energy transferred to each adatom is a more crucial parameter for the modification of many thin film properties, among them, its density. Thus, while at low ion energies only atom displacements at the surface should be expected, for higher ion energies penetration within the upper layers would occur and densification would be a near surface rather than a topmost surface process. These considerations are implied in the aforementioned model by Miiller (1986) and Netterfield et al., (1988), where densification takes place below the first monolayer according to a forward-relocation scheme. For certain materials such as c-BN or diamond it has been theoretically and experimentally determined that, for a given I/A ratio, the density of the films reaches a maximum value for a certain critical energy Ec. This energy is of the order of 100 eV and depends on the synthesised material. For these compounds, the evolution of density with ion energy follows a profile similar to that shown in Figure 3.11 (Reinke et al., 1997). It shows a sharp increase in densification that may be associated with the aforementioned phenomena of adatom mobility. However, well above the critical energy Ec, the impinging ions or the relocated atoms have enough energy so as to penetrate beneath the first surface layers (subplantation model, see section 5.12.3). This causes a significant fraction of the ion energy to be lost deeper in positions where the generated defects cannot be healed by newly arriving atoms. As a consequence, a progressive accumulation of lattice defects occurs (i.e., atom vacancies, point defects, atoms in interstitial positions, etc.), with the expected loss of crystallinity and density. This process is more significant for high ion energies and I/A ratios so that, for very high values, a severe amorphisation and decrease of density may occur.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
119
0.4
0.3
Q.
^ . 0.2
< 0.1
0.0
T
0
100
'
1
200
'
1
300
•
1
400
'
1
500
•
r
600
Ion Energy (eV) Figure 3.11. Densification of c-BN and caitoon as a function of the ion energy.
3.6. Defect generation The generation of structural defects (i.e., point defects, dislocations, etc.) in solids subjected to high-energy ion bombardment has been widely studied (Smidt, 1990 and ref. therein). In general, it is observed that the type and concentration of defects are very much dependent on bombardment conditions, such as type of ions and energy and fluency of the beam. Atom displacements and/or quenching of thermal spike regions are typical mechanisms leading to the generation of defects in solids. In some cases, ion bombardment may also contribute to the removal and healing of defects by supplying the energy required for lattice reconstruction.
3.6.1. Formation of defects in IAD thin films By contrast, in thin films the dependence of defect generation on bombardment and deposition conditions has not been so widely studied as in irradiated solids. The influence of ion beam induced defects in the nucleation processes during the initial stages of IAD thin film formation was discussed in section 3.2. For already grown thin films, a correlation has very often been found between macroscopic properties like conductivity, microhardness, optical properties, etc. and the amount of structural defects remaining in the films after preparation.
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Low ENERGY ION ASSISTED FILM GROWTH
The defects generated during IAD deposition can be of different types (e.g. vacancies, interstitials, point defects, clusters, dislocations, etc.). The previously mentioned amorphisation and the textural development phenomena that will be discussed in sections 3.7 and 3.9 can be considered as final situations resulting from the accumulation of atom displacement events. Under favourable circumstances it is possible to monitor intermediate situations where point or more extensive defects like dislocations are generated during IAD growth of thin films (Hultman et al., 1987). Thus, for TiN (100) epitaxial films grown by reactive magnetron sputtering with a bias applied to the substrate, the formation of dislocations loops has been observed by TEM. The concentration of this type of defect decreased as the substrate temperature and the bias voltage increased. Thus, the concentration of dislocation loops reached a minimum at intermediate voltage values to increase again for higher voltages. The increase of the energy of the ions must enhance the atom mobility, thereby accelerating the rate at which the defects can be annealed out during deposition. However, above a certain threshold value promotion of line defects is to be expected by the accumulation of point defects that, under such conditions, cannot be healed out at a sufficiently high rate.
3.6.2. Surface and bulk defects as a function of beam energy During deposition of IAD films, ion energy can be given up to the growing layer either at the surface or in the underlying regions. Atom displacements responsible for lattice damage are produced by energy deposition in bulk regions. By contrast, energy deposition at surface regions leads to atom displacements contributing to the packaging and smoothing of surfaces and interfaces (cf., section 3.3). Theoretical analyses have been carried out to determine, as a function of the ion beam energy and mass, the amount of energy deposited in the surface or underlying regions of the films (Ma et al., 2000). Figure 3.12 shows two plots for the energy per ion given up to the surface or the underlying bulk for N* and Ar+ ion bombardment of a series of carbon containing films. It is apparent from this figure that, for both surface or bulk atom mobilities, it is necessary to supply a given threshold energy (indicated in the figure only for surface driven processes). This threshold energy is higher for bulk than for surface mobility, thus resulting in an energy window where it would be possible to enhance surface mobility without affecting the atom positions of subsurface layers. The existence of such an energy window can be of the utmost importance for the preparation of epitaxial thin films where the atom positions of substrate should not be altered (cf., section 3.3.5). Above the threshold of bulk
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
121
atomic mobility, it should be expected that the number of atom displacements induced in the bulk increases and mat, therefore, a more defective thin film is synthesised under these conditions. However, it must be noted that a counterbalance effect of this tendency that consists of the removal of already formed defects by the mobilised atoms is not taken into account in Figure 3.12. 10J ;
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Figure 3.12. Deposited energy producing atom mobilisation in surface and underlying bulk vs. incident energy and mass of projectiles into graphite, diamond and silicon carbide: (a) for nitrogen ion; (b) for Ar ion. Reproduced from Ma et al. (2000) with permission.
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Low ENERGY ION ASSISTED FILM GROWTH
3.6.3. Defects and control of the microstructure of thin films by annealing treatments Residual defect concentration in IAD thin films may have great importance for an effective control of the microstructure of the film by post-deposition annealing treatments. Changes in mechanical, electrical and other thin film properties may be a result of these treatments, whereby the concentration and type of structural defects present in the film may be of primary importance for the control of these and other properties. Thus, a detailed knowledge of the defects present in the films is usually necessary to define the best annealing conditions required to get thin films with defined properties. Since the activation energies for diffusion and, in general, the thermal stability of each type of defect are different, a precise knowledge of these defect characteristics is needed for a tailored processing of thin film properties. A common technique to monitor defects in materials is X-ray diffraction (XRD). In general, broad diffraction peaks are an indication of a defect-rich microstructure. Such a broadening can be due to either size effects (i.e., small size of crystalline domains) or to an inhomogeneous distribution of microstrains. Both situations can be differentiated by measuring two or more order peaks of a given diffraction plane (e.g. (200) and (400), etc.) and comparing their respective broadenings. For vacuum arc deposited CrN coatings prepared by applying different substrate voltages, it has been shown that XRD broadening is mainly due to inhomogeneous strain distribution within the film (Aimer et al., 2000). Figure 3.13 shows the evolution of XRD peaks for samples prepared by applying different substrate voltages. For both reported films, the systematic decrease in peak width has been interpreted as being due to the release of inhomogeneous distribution of strain within the films, while the shift in peak position has been attributed to the accumulation of residual stress. Annealing at increasing temperatures produces a decrease in peak intensity that has been interpreted as a release of this inhomogenous strain distribution. Thus, it is interesting to see in Figure 3.13 that at intermediate temperatures the variation of the peak width for the 50 V sample is more pronounced than for the 300 V one. From a careful analysis of the evolution with temperature and time of the peak widths, it is possible to deduce the value of the apparent activation energies for the annealing of defects, a parameter whose knowledge can be critical for an effective control of the microstructure of the CrN films after annealing.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
123
V=50V
Cr2N(300)
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Figure 3.13. X-ray diffraction patterns as a function of tempering temperature for (a) Vs=50 and (b) 300 V CrN coatings. Dotted lines indicate equilibrium 290 positions and arrows indicate approximate 28 positions in the as-deposited coatings. Small amounts of Cr exist in the original film, while the Cr2N phase is formed as a result of the annealing treatments. Reproduced from Aimer et al. (2000) with permission.
3.6.4. Inert gas incorporation The incorporation of inert gases in thin films is a well-known phenomenon that has been recognised and discussed for many years (Carter et al., 1980). Incorporation of inert gas atoms within a film can be considered as equivalent to the introduction of a specific type of structural or microstructural defect. Experimentally, it has been found that thin film properties like electrical resistivity or microhardness, are dependent on the amount and type of incorporated inert gas (Paturand et al., 1999).
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The amount of incorporated inert gas atoms may reach up to five or more atomic percents in thin films grown at room temperature by IAD. Increasing the substrate temperature generally contributes to the release of the inert gas, so that the incorporated amount may drastically decrease at high temperatures. Since the ions need to penetrate within the bulk to become incorporated in the film, the probability of inert gas incorporation depends on the ion energy and mass. Trapping probabilities have been theoretically estimated and experimentally determined for a large variety of inert gas and target materials. It is found that the ions need a threshold energy to be effectively trapped and that above this energy the incorporation rate sharply increases with the ion energy. It is also shown that low mass atoms become more effectively trapped than heavier atoms. This trend has been justified because the reflection probability of the inert gas ions decreases for the lighter elements (Eckstein et al., 1986). Incorporation of inert gas atoms in thin films can be proven by RBS, XPS or any other atom sensitive technique (Holgado et al., 2001). Figure 3.14 shows some RBS spectra corresponding to Zr0 2 thin films prepared by IBICVD using either 0 2 + or 0 2 + + Ar+ ions. In the spectrum of the second sample, an additional structure, shown in an expanded scale in the inset, indicates that Ar (about 5% atomic) is homogeneously distributed within the film. In this experiment, it was also found that the incorporation of Ar in Zr0 2 films provokes the stabilisation of the tetragonal structure of this oxide. Stabilisation of this structure at room temperature is generally achieved by doping with other cations like Y3+, Ca2+, etc. Although the influence of the incorporated inert gases in modifying certain thin film properties has been recognised in many papers, the possibility that incorporated Ar may produce changes in the crystallographic structure of oxide or nitride thin films (e.g. crystalline phase, plane orientation, amorphisation, etc.) is an interesting feature which deserves future investigation. A common effect of the incorporation of inert gas atoms in the structure of a thin film is an increase in the compressive stress (Fang et al., 1993). This point will be dealt with in detail in section 3.10.6. Here, it can be noted that films with high compressive stress may have unexpected structural properties and that the accumulation of stress can be a reason for the stabilisation of metastable crystalline phases as in the previous example of Zr0 2 .
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
125
Zr
Ar
MNx II
Ar ) - Zr0 2
(0 2 *+ Ar*) - ZrO ; "As Prep."
J
100
200
300
400
500
Channel Number
Figure 3.14. Experimental and simulated RBS spectra of Zr0 2 thin films deposited on Si by IBICVD by using 0 2 + or mixtures 0 2 + + Ar+. The inset shows, in an enlarged scale, the region corresponding to the embedded Ar that remains in the films even after annealing at 773 K. Reproduced from Holgado et al. (2001) with permission.
3.7. Amorphisation, crystallinity and phase transformations In previous sections, we have discussed the fact that ion bombardment during thin film growth produces significant modifications in the crystallographic structure of the thin films. The development of specific textures (i.e., preferential orientation of some crystallographic planes) (cf., section 3.9) or the accumulation of structural defects (cf., section 3.6) is some important lattice transformations that are or will be discussed in this chapter. In the present section, we would like to comment on other crystallographic effects that can occur in IAD films. These effects refer to amorphisation phenomena and the production of metastable phases. Regarding this second aspect, a more detailed account will be given when dealing with the synthesis of c-BN, diamond and other related materials prepared by using IAD techniques (cf., Chapter 5). Here, we will comment briefly about amorphisation and on the ion-beam stabilisation of out-of-equilibrium phases observed in some oxide thin films.
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Low ENERGY ION ASSISTED FILM GROWTH
3.7.1. Amorphisation in IAD thin films The term amorphisation is not very precise and a clear definition of this concept will require a deep discussion that is outside the scope of this book. Here, we will use the intuitive idea that amorphisation refers to the lack of crystal order on a large scale. Such a situation can be proved experimentally by XRD or other techniques of crystal structural analysis (e.g., when the XRD diagrams do not present well defined peaks). Note that, although the order at large scales monitored by these techniques may be missing, local order may still exist. Structural amorphisation of thin films may occur by rapid quenching of the atoms, positions in thermal spike zones or similar damaged areas of the thin films (cf., section 1.6.2) or by an excessive accumulation of defects. With the adopted definition, relying on negative experimental evidence by XRD, it is even possible to consider amorphisation as resulting from too small a size of crystalline domains, a situation that would lead to very broad peaks in the XRD diagrams. Amorphous structures in IAD films are more easily obtained in compounds with ionic and covalent character (i.e., oxides, nitrides, diamond and related compounds, etc.) but very seldom with metals, where the atoms usually have enough mobility to occupy their correct positions in a crystal lattice.
3.7.2. Effect of temperature on crystallisation In many cases, partial rather than complete amorphisation can be found in IAD thin films. Experimental evidence of partial amorphisation can be inferred, for example, when broad structures instead of flat backgrounds with no peaks appear in the XRD diagrams, when the Raman peaks of structural modes of vibration become broad or when the electron diffraction diagrams present diffuse rings. Figure 3.15 shows as an example the XRD diagrams of thermally grown Ce0 2 films prepared by ion beam sputtering of a target and deposition on a substrate kept at different temperatures from 50° to 500°C (Karakaraju et al., 1997). The films were not ion beam assisted during growth. Partial amorphisation can be deduced from the stepping background appearing in the diagrams between 40° and 25° for the samples prepared between 50 °C and 300 °C. The large width of the peaks within this temperature interval constitutes a clear indication of a low crystal order, likely because of a small size of the crystal domains. In this experiment it is also
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
127
interesting that, as deduced from the systematic decrease in peak intensity observed in the XRD diagrams, progress in crystalline disorder occurs from 50° to 300°C. Then, suddenly, the film becomes fully crystalline at 400 °C, as deduced from the very narrow peaks recorded by XRD, while it also develops a (200) preferential texture. Texturing of Ce0 2 thin films will be discussed in more detail in section 3.9.3 for IBAD Ce0 2 thin films whose growth is assisted with ions of 100 and 300 eV supplied by an independent ion source. In that case, the films are very crystalline for most of the deposition conditions, probably because of the enhanced atom mobility induced by the ion bombardment. In the present experiment, however, fine grains with a high degree of disorder grow at T < 300 °C. Above this temperature, thermal annealing effects seem to be more important and the films become crystalline and preferentially oriented.
(Ill)
Po2 = 0.01Pa (220)
(200)
1
1 I
^..
I
S i (222) Powder
..1
(311) 500°C
*v?\3!
400°C (311)
(220)
300°C
200°C
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50°C .
25
30
35
40
i
i
45
50
.
1
55
.
1
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26 Figure 3.15. X-ray diffraction diagrams of ceria films deposited at different substrate temperatures by ion beam assisted deposition by bombardment with Ar+ ions of 1 keV energy. Reproduced from Karakaraju et al. (1997) with permission.
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Low ENERGY ION ASSISTED FILM GROWTH
3.7.3. Amorphisation and phase transformation phenomena. Stabilisation of unstable phases Amorphisation may occur in oxide thin films as an intermediate stage appearing when a given crystallographic phase of an oxide thin film is being transformed into another one. Such phenomena have been reported for example for Ti0 2 thin films prepared at room temperature by FVAD with the substrate biased up to -400 V. Ti0 2 present several crystalline phases. For the bulk material, anatase is the stable phase at T<600°C, while rutile stabilises at T>600°C. A well-crystallised anatase structure is obtained for the FVAD thin films when the substrate was not polarised. Then, for polarisation voltages below -100 V, amorphous Ti0 2 (i.e., no XRD diffraction peak pattern) was obtained, while for higher polarisation voltages the rutile structure was formed with enhanced crystallinity as the absolute value of the substrate voltage increased (Bendavid et al., 1999). In other IBICVD experiments, stabilisation of the rutile phase has been reported for thin films grown under the assistance of an external ion beam of 10 keV energy (Caballero et al., 1995) even if the substrate is not externally heated during deposition. The fact that rutile can be formed by ion bombardment at temperatures lower than 600 °C has to be related to the large amount of energy and momentum transferred to the growing films by the impinging ions. Stabilisation of high temperature phases at low substrate temperatures has been reported for other oxide systems such as ZnO or A1203 prepared by IBAD methods (Schneider et al., 1997). In the latter case, formation of the phases 0 and K, whose development usually requires very high temperatures, has been reported for temperatures as low as 400 °C.
3.7.4. Monitoring the degree of amorphisation in IAD thin films One way of estimating the degree of crystallinity of a thin film is by looking to the width of the rocking curves of a given diffraction peak. Figure 3.16 shows a plot where the width Aco of the rocking curve of the (200) diffraction peak of an YSZ thin film grown by DIBS by Ar+ or 02++Ar+ ion bombardment is compared with the normalised energy of the beam (En=Eion (I/A), section 1.8.2). This parameter can be interpreted as the average energy delivered per deposited atom during the growth of the film. The results reported in this figure demonstrate that for the films subjected to bombardment with low E„ of the order of 50-60 eV, Aco presents the maximum
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
129
values, probably because at such low En values insufficient energy is provided to the growing films as to significantly improve their crystallinity. For values of En around 100 eV, A(0 has a minimum value and, therefore, the films experience a significant improvement in their crystal order. However, when higher E„ values are used an increase in Aco is again found. This increase must be associated with a loss in crystallinity, likely resulting from the generation under such conditions of a higher concentration of defects (Koch et al., 1997). On the other hand, the different shape of the curves in Figure 3.16, corresponding to families of thin films prepared either by Ar+ or Ar + +0 2 + bombardment, can be related to the higher mass of Ar atoms (after impinging on the surface, the 0 2 + molecular ion dissociate into atomic species) and its higher efficiency for the preferential sputtering of oxygen. Under these conditions, it is expected that the Zr mobility through the oxygen defective lattice increases, thus enabling the cation sublattice to become more ordered, as it is found experimentally.
Mill
1
I I I Mil]
I
rmj
1
i i (IITI|
r—TTTTT
— O D O A V • • A T
3.5° 3.0° S 2.5° < 2.0° 1.5° 1.0°
l|l||
10°
I 1 I 1J1111
r T 11 Mill
I
101 102 Energy / Atom
I J JUIll
guide - to - the - eye 50 eV Ar* 60 eV Ar* 120eV Ar* 180 eV Ar* 240 eV Ar* 60 eV Ar*/02* 120eV Ar*/02* 180eV Ar*/02* 350 eV Ar*/02*
I J I mil
103 (eV)
104
Figure 3.16. Widths of the XRD rocking curves, Am (FWHM) of YSZ films prepared under different IBAD conditions at 580 °C as a function of the normalised energy E„. Open symbols, IBAD results obtained with an Ar* ion beam; closed symbols, IBAD results obtained with an Ar70 2 * ion beam. The various symbols represent experiments with different ion energies as given in the inset. Reproduced from Koch et al. (1997) with permission.
An analysis of the crystallinity of a thin film relying exclusively on the interpretation of XRD diagrams faces the problem of estimating quantitatively the
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Low ENERGY ION ASSISTED FILM GROWTH
degree of amorphisation. This is, in fact, not possible because this technique is blind to amorphous structures. A semi-quantitative estimation of the amorphisation degree of a thin film can be done by combining XRD with an X-ray absorption spectroscopy (XAS) investigation of the structure of the film (Caballero et al., 1995). This technique is sensitive to the local environment around a given atom and, therefore, provides average information about the coordination spheres of all the cations in the sample, independently of the fact that this environment corresponds to a crystalline or amorphous phase. Figure 3.17 shows the Ti K-edge X-ray Absorption Near Edge structure (XANES) of partially crystalline thin films of Ti0 2 prepared by IBICVD under different conditions of deposition. Reference XANES spectra of a completely amorphous thin film is included for comparison. The spectra of the thin films can be reproduced by linear combination of the spectra of anatase, rutile and amorphous phases. This linear combination includes a significant contribution of the amorphous reference spectrum, indicating that there is a high degree of amorphisation that can vary from 10 to 95% depending on the conditions of preparation. It must be stressed that the XRD patterns of these samples showed peaks of anatase and/or rutile and that the amorphous contribution to the structure is not detected by the conventional XRD analysis. The study can be complemented by looking at the extended X-ray absorption fine structure (EXAFS) of the XAS spectrum which provides a clear description of the average coordination state of titanium (i.e., average coordination number and Ti-O distances). The percentages of the anatase, rutile and amorphous phases in the different samples reported in Figure 3.17 supports the concepts explained in this section about the dependence of amorphisation on experimental conditions (i.e., ion energy) and the development of amorphous phases as intermediate stages in the transformation of one crystalline phase into another under ion bombardment during thin film growth (cf., section 3.8.2).
3.8. Compound formation by IAD If, instead of inert gas ions, a reactive ion beam is used to bombard the film during its growth, a new compound or phase can be obtained. Oxides and nitrides are the most typical materials prepared by using this IAD approach. For the synthesis of thin films of these materials, bombardment with accelerated 0 2 + or N2+ ions or mixtures of these ions with Ar+ or other inert gas ions are used. Control of the stoichiometry and the depth distribution profile of thin film components or the
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
131
preparation of new metastable phases are some of the most interesting applications of using reactive ion beam assisted deposition procedures. 1kV(300K)
4940 4960 4980 5000 5020 5040 5060
energy / eV Figure 3.17. XANES analysis of the contribution amorphous and crystalline phases of Ti02 in thin films prepared by BICVD according to different protocols as indicated. Thick line: experimental spectra. Thin line: calculated spectra. A reference spectrum of amorphous Ti02 is included in the figure for comparison. Reproduced from Caballero et al. (1995) with permission.
3.8.1. Control of stoichiometry in IAD thin films Control of the stoichiometry of the thin films is possible by adjusting the I/A ratio during deposition. Classical examples of this stoichiometry control can be found in literature for many nitride (e.g. A1NX, SiyNx or TiNx sub-stoichiometric phases)
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Low ENERGY ION ASSISTED FILM GROWTH
(Hirvonen, 1991 and references therein) and oxide compounds (e.g. CuOx, PbOXi etc., (Hirvonen, 1991 and refs. therein; Guarnieri et al., 1989). In all cases it is found that the amount of reactive gas atoms incorporated in the film increases with the I/A ratio, and reaches saturation for a target composition equivalent to that of the thermodynamically stable phase (e.g. A1N or Si3N4). For aluminium or silicon nitrides, it has also been shown that the amount of incorporated nitrogen (i.e., ratio between the nitrogen and metal atoms in the film) follows an almost linear relationship with respect to the I/A ratio with a slope close to one. In general, above the stoichiometric limit no more N or O can be incorporated within the thin film structure and the excess of nitrogen or oxygen is released into the deposition chamber. Another factor contributing to the effectiveness of the incorporation of either nitrogen or oxygen to the film is the ion energy. Higher values of this parameter usually stimulate the incorporation processes of either oxygen or nitrogen into the film. Using mixtures of reactive and inert gas ions for assisting the deposition can produce a similar effect. The addition of inert gases aims to produce more favourable momentum transfer conditions owing to the higher mass of Ar or other inert gases with respect to atomic oxygen or nitrogen (cf., section 1.8.2) (0 2 + or N2+ molecular ions dissociate when impinging on the surface of the target). In this way, an increase in adatom mobility is produced that can lead to an additional incorporation of reactive atoms within the thin film structure. For many compounds, the amount of incorporated nitrogen or oxygen in the film also depends on the partial pressure of these gases in the reaction chamber. Fresh surfaces of metal films like Ti, Al or Si are very reactive towards N 2 or 0 2 molecules and may react with them, even at low temperatures, in the absence of ion bombardment. The extent of the chemical reaction depends on the partial pressure of the gas in the chamber. Ion bombardment, besides its beneficial effects in terms of densification, microstructural control, etc., strongly enhances the incorporation rate of reactive atom species. Accelerated ions supply some additional energy to the growing film that promotes processes such as implantation or atom mixing and mobilisation that may favour the incorporation of extra amounts of oxygen or nitrogen atoms within the network. A comparison of the effectiveness of nitrogen incorporation into a growing titanium nitride thin film with and without ion bombardment is shown in Figure 2.4.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
133
Owing to the direct relationship between the amount of oxygen or nitrogen incorporated within the film and the I/A ratio and beam energy, the fabrication of thin films with a graded composition may be very straightforward by means of IAD methods. Graded A1NX, SiOx, SiNx or other oxide and nitride thin films can be very interesting for many technological applications and the graded synthesis of some of these compounds has been reported in the literature (Vechten et al., 1990; Watanabe et al., 1999) (cf., rugate filters, section 4.4.4). Besides parameters like ion energy or I/A ratio, it is known that the charged or neutral character of the accelerated species impinging on the target may influence the thin film properties. Accelerated N2 species (i.e., accelerated neutral molecules) or N2+ ions have been used to grow GaN thin films by evaporation of Ga and simultaneous N2 or N2+ bombardment. GaN, together with A1N and InN, are promising materials for optoelectronic and high temperature electronic devices. The GaN thin films prepared by bombardment with neutral N 2 species have a different texture, higher crystallinity and density and lower surface and interface roughness than samples prepared with N2+ ions of a similar kinetic energy. The photoluminescence properties of the films were also superior as expected for thin films with less residual stress and lower concentration of defects (Kim et al., 2000).
3.8.2. Metastable phases of nitride thin films An interesting application of reactive IAD is to obtain crystalline metastable phases of certain compounds with unusual compositions. Generally, these phases cannot be prepared by chemical vapour deposition (CVD) or other thermal methods. The stabilisation of different crystal phases of BN or carbon films will be discussed when dealing with these thin film compounds in Chapter 5. Another widely studied material has been titanium nitride. The Ti-N phase diagram shows that at room temperature it is possible to stabilise three different phases with different nitrogen content: a-Ti, e-Ti2N and 5-TiN. The e-Ti2N is only allowed in a very limited compositional range and its synthesis in the form of thin films is rather difficult by procedures other than reactive IBAD. In fact, the synthesis of pure e-Ti2N is possible by IAD methods by using high energy N2+ ion species. The procedure used has sometimes been named ion beam mixing because of the similarities existing with this procedure of preparing new compounds (cf, section 3.4.1). By properly adjusting the ion beam current and evaporation ratio of Ti (i.e., I/A ratio) it is possible to obtain either the conventional 8-TiN phase or the more unusual e-Ti2N
134
Low ENERGY ION ASSISTED FILM GROWTH
while, for intermediate conditions, mixtures Ti2N + TIN or Ti + Ti2N + TiN are obtained (Kiuchi, 1993). The need for high energy ions for obtaining such a new compound would indicate that it is formed in deep regions of the film where a kind of ballistic mixing induced by the penetrating N1" ions is the factor controlling the formation of the e-Ti2N phase. Chromium nitride thin films are very interesting materials used for anticorrosion applications. CrN and Cr2N are typical stoichiometric phases of chromium nitrides. The synthesis of chromium nitride thin films, and in particular CrN, is more difficult than that of titanium nitrides. This is due to the very narrow region of stability of CrN in the phase diagram and, in general, to the reduced thermodynamic stability of chromium nitride and the lower chemical activity of chromium towards nitrogen in comparison with that of titanium. Chromium nitride thin films can be prepared by IBAD by bombardment with rather energetic (some tens of keV) N2+ ions a growing Cr film deposited by evaporation, whereby the control of the overall stoichiometry of the system is achieved by adjusting the I/A ratio. Figure 3.18 shows the regions of existence of defined phases or mixtures of phases of the Cr-N system. Composition, and, therefore phase stabilisation of chromium nitride IBAD thin films, are solely determined by the I/A ratio of the impinging N2+ species since no incorporation of nitrogen occurs by reaction of this gas with a fresh chromium surface. In Figure 3.18, it is apparent that while the regions of existence of the nitrogen containing chromium (i.e., Cr in the plot) and Cr2N are wide, it is less well defined for CrN, probably as a result of the fact that more nitrogen has to be delivered to the system to reach such a stoichiometry. Owing to the small window of existence of pure CrN, the actual films are a mixture of crystallites of different phases, with a distribution that, as indicated in the figure, depends on the experimental conditions of deposition (Ensinger and Kiuchi, 1997; Volz, etal., 1998). The mechanism by which these unusual titanium or chromium nitride phases are formed during high energy IAD deposition of nitride thin films has been considered to be similar to an ion implantation process in a bulk substrate. In fact, since the impinging ions have a very high energy they are only stopped at a relatively large distance from the surface of the growing thin film. It is at these deep regions rather than at the surface where the excess nitrogen required to get a CrN composition becomes incorporated after the ions lose their final energy and where the formation of a crystalline CrN or e-Ti2N phases might eventually take place.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
135
Note that the release of excess nitrogen would be hindered from these deep regions, thus favouring the formation of these nitrogen rich phases.
40
80
120
160
200 2
ion current density [jlA/cm ] Figure 3.18. Regions of existence of different phases in a deposition rate/ion irradiation intensity diagram. "O" denotes that there is not net film growth due to extensive sputtering of the film. Reproduced from Ensinger and Kiuchi (1997) with permission.
Another interesting example of the preparation of several nitride phases with different compositions is that of zirconium nitride. While ZrN has a metallic character, Zr3N4 is an insulator. Both phases can be obtained by DIBS at relatively low ion energies by changing adequately the experimental conditions of preparation. Zr3N4 can be obtained at low temperature and N2+ beam energies around 200 eV. However, higher temperatures and lower kinetic energies of the ion beam yield a ZrN phase. (Pichon et al., 1999). The mechanism controlling the final thin film composition in Zr3N4 seems to involve the incorporation of the excess nitrogen in subsurface regions. Subsurface implantation occurs for ions with energy above a certain threshold value (cf., section 5.12.3), but not for energies below that limit. It is interesting to note the similarity of the mechanisms involved in the synthesis of nitrogen-rich phases either by using high energy (i.e., Ti and chromium nitrides) or low energy ions (i.e., zirconium nitride), which in both cases consists of sub-surface processes.
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ZrN and Zr3N4 have quite different properties. For example, while ZrN is a good conductor of electricity, Zr3N4 is an insulator. Figure 3.19 illustrates how ion bombardment may transform the electronic properties of the compounds by changing its stoichiometry (Sanz et al., 1998). This figure shows a series of photoemission spectra from the valence band region of Zr subjected to ion bombardment. Incorporation of increasing amounts of nitrogen leads to a decrease in the concentration of electronic states at the Fermi level (i.e., binding energy zero) and to the evolution of a big peak at a binding energy around 5 eV. For a Zr/N ratio equal 1.06 (i.e., ZrN) there is still a certain concentration of electronic states at the Fermi level, thus confirming the metallic character of this material. However, for a Zr/N ratio equal to 1.33 (i.e., Zr3N4), no electronic states appear at this position in agreement with the insulating character of this compound.
15
10
5
0
BINDING ENERGY (eV)
Figure 3.19. Photoemission spectra of the valence band region of Zr subjected to successive doses of ion bombardment. Formation of ZrN and ZrsNi phases is deduced from the stoichiometric ratio between Zr and N. Reproduced from Sanz et al. (1998) with permission.
The previous examples show that by modifying the experimental conditions it is possible to prepare thin films of unusual phases with controlled compositions. However, much experimental work is still required to, firstly, investigate the possibility of preparing new metastable phases of other materials and, secondly, to define a general model to account for the mechanisms that lead to the formation of such phases.
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3.9. Texture development Texturing is a very common phenomenon for many types of polycrystalline thin films prepared by different methods. This implies that the elemental crystallites that constitute the films are not randomly oriented, but distributed in such a way that they present a common preferential orientation for given families of crystal planes. Generally, this preferential orientation is referred towards the thin film surface. A very interesting feature for IAD thin films is the possibility of modifying the texture by the deposition conditions (i.e., ion energy, I/A ratio, etc.). In these thin films, the texturing degree according to certain planes can be very high (sometimes close to 100%) and the preferential orientation of one or another type of crystal planes can be conveniently tailored according to particular thin film applications. The degree of understanding of the induction mechanisms of these processes by ion beam bombardment is relatively high and there are several models that explain reasonably well these preferential orientation effects. This section deals with the phenomenology of the texturing process, the models proposed to explain these effects and the technological implications of controlling the texture of IAD films.
3.9.1. Monitoring the texture in IAD thin films by XRD: basic definitions The development of a preferential orientation of crystal planes in thin films can be easily recognised by electron or X-ray diffraction (XRD) techniques. In the following, we will present some basic definitions and concepts that are usually applied to investigate preferential orientations in thin films. A more detailed description of the XRD techniques and methods used for these applications can be found in more dedicated publications (Barrett et al., 1980). XRD is the most popular technique for texture characterisation of thin films. With this aim the diagrams have to be recorded with a 6-26 configuration (i.e., Bragg-Brentano configuration). Figure 3.20 shows schematically the disposal of an x-ray source and a detection system with respect to the plane of the thin film for this type of analysis. In this configuration, while the sample rotates an angle 6, the detector rotates by an angle 26. Under these experimental conditions, only planes parallel to the thin film surface can be detected, because they are the only ones that can fulfil Bragg's law. In fact, to detect a family of crystal planes (hkl) it is
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a requirement that the line perpendicular to them and the incident and reflected xray beams are on the same plane, and that the equation 2d Sen 0 = n A holds. In this equation, d is the distance between the planes of the (hkl) family and 0 is the experimental diffraction angle according to Figure 3.20. Planes that are not parallel to the thin film surface are not detected in this experiment.
X-ray source
Figure 3.2§. Schematic description of the experimental arrangements to characterise textures in thin films. 8/20 scan angles; CD angle to measure rocking curves; y aamuthal turning angle to measure polar plots.
As an example, Figure 3.21 shows idealised XRD. diagrams of an fee metal thin film with a randomly oriented distribution of crystal planes and that of a thin film of the same material that has developed a preferential (100) texture. The preferential growth of this latter family of planes with an orientation parallel to the thin film surface is easily recognised in the diagram by the relatively higher intensity of the (200) diffraction peak as compared with the randomly distributed sample. The texturing degree of a thin film (i.e., the extent by which their crystallites grow with certain planes oriented parallel to the surface) can be quantitatively defined by using the so-called texture coefficient of a given plane family (hkl). This coefficient is defined as:
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l
hkl
-
'A,U
hkl 5 hkl
139
(3.4)
'A,U
where Iha is the measured intensity of the (hkl) peak in the film, fhkt is the intensity of the same peak for a randomly oriented sample and n is the number of peaks considered in the analysis. A value of T equal to one means that the grains are randomly oriented, while greater than one means that the film presents a preferred orientation of that family of planes. In the idealised diffraction pattern of Figure 3.21, the texturing parameter of the (200) planes is 3.65.
(200) (111) 3
J
3. 3> d)
_c
(220)
(311)
1
A
(111)
r
I (200) 1(220)
i
30
......
40
.
1
1
50
60
•
1
—
70
•
i
•
80
Figure 3.21. Idealised X-ray diffraction patterns of a randomly oriented (bottom) and a (100) textured (top) fee metal thin film.
To simplify the calculation, it is also very common to use a coefficient of the degree of orientation R that is defined as:
R =
1
hkl
(3.5)
It
2~i Wi i=i
In this case, no reference of a random material is required. R is always less than one except for thin films where only a given peak (hkl) is detected in the
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diagram. Such a situation would mean that the film is completely oriented according to that family of planes. An interesting situation is when a family of preferentially oriented planes contains or is perpendicular to a given crystallographic axis. This means that such crystal axes would be either parallel or perpendicular to the thin film surface (e.g., in the (100) textured fee thin film of Figure 3.21 the crystallographic c axis is preferentially oriented parallel to the surface of the sample). To get a more detailed description of the texture of the thin films, the 9/26 scans can be complemented with the recording of the so-called "rocking curves" and "polar plots". According to Figure 3.21, the (200) planes are preferentially oriented parallel to the thin film surface. What usually happens is that besides the series of completely parallel planes detected with the Bragg-Brentano geometry, there are other planes of the same family that forms a small angle ft} with respect to the sample surface. These planes can be detected if the sample is tilted at an angle ft} as indicated in the scheme of Figure 3.22.
Figure 3.22. Arrangement of a family of (hkl) planes parallel to the thin film surface (middle) and forming a certain angle coi with respect to that orientation.
A practical way of getting information about the deviation in the orientation degree of the planes with respect to the sample surface is by recording a "rocking curve". In practice, it is recorded by selecting the angle 0 of the maximum of a diffraction peak and then by tilting the sample by an angle ±co around the horizontal position. The obtained curve resulting from a representation of the intensity of the peaks with respect to the magnitude of the angle ±co is called a "rocking curve" and its width gives information about the degree of orientation of
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this family of planes with respect to the surface of the film. Very sharp rocking curves indicate a high degree of order (i.e., most crystallographic planes of that family are parallel to the surface of the thin film). However, broad rocking curves are proof of the existence of a great number of (hkl) planes whose orientation deviates with respect to the plane defined by the surface of the film. Besides the detection of the families of planes which are oriented parallel to the surface, it can be interesting to get some evidence of the distribution of planes oriented according to other directions and, in particular, whether the crystallites present a biaxial orientation. From an experimental point of view, this situation is similar to that existing in an oriented single crystal, with the obvious difference that pollycrystalline films are composed of many elemental crystallites. It is said that the film has a biaxial orientation when two crystallographic directions of all, or the majority of the film crystallites, have the same orientation. If two crystallographic axes are oriented parallel to the thin film surface, then it is referred to as "in plane " biaxially oriented thin films (cf., Figure 3.29). Biaxial orientation in thin films can be ascertained by recording the so-called "polar plots" of a given family of planes. Recording of a polar plot is carried out by following the intensity of a given diffraction peak as the sample is turned around its azimuthal axis 7 at giving values of the tilting angle (cf., Figure 3.21). With this type of polar plot it is possible to establish whether there are crystallites with their crystallographic planes preferentially oriented according to certain directions with respect to the normal of the film surface. A polar plot is equivalent to a stereographic projection of the crystal direction of a chosen family of planes (h'kT) on another one (hkl). For the analysis of thin films, this latter is generally taken as the planes that are oriented parallel to the sample surface. In a "polar plot", peak intensities are projected on a circular plot according to a cartographic procedure as a function of the angles (j> (or co, see Figure 3.20) and y. An example of such types of polar plots is represented schematically in Figure 3.23 for a diffraction peak of a film where the crystallites are single oriented towards the thin film surface (left) and for a film with biaxially aligned crystallites. In the first case, the polar plot of a given family of planes (h' k' 1') forming a certain angle with respect to the surface but randomly oriented in the other directions renders a ring in the plot. Meanwhile, a biaxially oriented film furnishes four zones of maximum intensity. Each zone is produced by families of planes related by symmetry (e.g., in a cubic system the (1 0 0) planes include the families
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(1 0 0), (0 1 0), ( 1 0 0) and (0 1 0)). For a perfectly oriented single crystal, such a plot would convert into four points. If the family of planes (h1cT) in the different crystallites were completely random oriented, the polar plot would not give any zone or ring of maximum intensity. The degree of biaxial ordering in a real thin film is higher as these four zones are less spread in the stereographic projection. Spreading of the intensity area would reflect a situation similar to that existing when the rocking curve is broad (i.e., when there is some deviation of the crystal planes parallel to the surface with respect to a perfect horizontal orientation), although in this case, the referred crystal directions form a certain angle with respect to the normal of the thin film surface.
Figure 3.23. Schematic representation of two polar plots of a family of crystallographic planes (hlcT) that form a certain angle with respect to the thin film surface but are randomly oriented in other directions (left) or present a biaxial orientation (right).
3.9.2. Texture in PVD thin films When the substrate temperature during growth is high enough to induce high adatom mobility, PVD thin films normally grow with the most closely packed planes parallel to the thin film surface. Thus, face centred cubic (fee) metals develop a (111) texture, body centred cubic (bec) metals a (110) texture and the hexagonal compact package (hep) metals a (0002) texture. The reason for a preferential orientation of these planes parallel to the surface relies on energetic factors: the adatoms tend to minimise the energy of the system by ordering and aggregating in those crystal planes that have a minimum free energy. However, the extent of the
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preferential orientation is rather limited and very much dependent on kinetic restrictions related to the temperature of the substrate, shadowing effects, arrival rate of evaporated material, etc.
3.9.3. Texture evolution in IAD thin films and process parameters Preferential orientation can be enhanced or modified with respect to the usual situation in PVD films by assisting the growth with ion beams. The evolution of a preferential texturing has been observed for a large variety of thin film materials including metals, oxides, nitrides, etc. Texturing can occur according to one direction (i.e., with respect to the direction normal to the surface of the film). For some compounds and under certain experimental conditions, texturing may also occur for two biaxial "in-plane" directions (i.e., two crystal axis are preferentially oriented with respect to the thin film surface). The degree of preferential orientation is very much dependent on process parameters such as the ion energy or the I/A ratio. Usually, an enhancement of the preferential orientation is achieved as the magnitude of these two parameters increases. In metal thin films, a clear dependence is always found between texture development and ion energy. In this case, it has been argued that one of the factors controlling the preferential growth of a given family of planes is the minimum free energy of the system. According to this criterion, fee metal films tend to develop (111) preferred orientations. However, the development of alternative preferential orientations is also possible, being very much dependent on the ion beam energies and on the efficiency for energy transfer between the impinging ions and the target atoms. For Ni films grown under ion bombardment, it has been shown that the tendency to change the preferential orientation from a (111) to (200) and (220) textures, the latter two with higher surface energies than the former, follows the amount of electronic energy that is transferred by ion bombardment. In an atom collision, the effective nuclear (Sn) and electronic (Se) transferred energies depend on the mass of the two atoms involved in the collision and on the ion energy (cf., section 1.3.1). The energy transferred through these two types of interactions can be calculated analytically or by using the TRIM code (cf., section 1.4). For Ni it has been shown that the Se/Sn ratio increases when the beam energy grows from 400 eV to 10 keV. The aforementioned evolution from a (111) preferential orientation to another characterised by (200) and (220) preferential orientations follows the same trend. Theoretical considerations of ion-atom collisions reveal that in metals the
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energy transferred through electronic energy losses are easily transformed into thermal energy of the target atoms. Therefore, it is expected that as the Se/Sn ratio increases, the thermalisation of atoms is more favourable and Ni may develop a texture with a higher internal energy associated with less favourable plane orientations (Kuratani et al., 1997). Texturing is also a common effect in IAD oxide thin films. Figure 3.24 shows the 9-28 XRD diagrams of a series of Ce0 2 IBAD films, where it is possible to follow the evolution of the texture of this oxide characterised by developing preferential orientation when subjected to ion bombardment during growth as a thin film. These thin films have been grown by magnetron IBAD by assisting the deposition with ions of energies between 100-300 eV that form an angle of 55° with respect to the substrate normal. It is apparent in the figure that the films become more textured according to the [002] as the ion energy increases. This means that the c crystallographic axis becomes preferentially oriented in a direction perpendicular to the film surface. A high degree of preferential orientation is found for the 300eV film where only the (002) plane diffraction peak is visible in the diagram (Gnanarajan et al., 1999). The progression of the preferential orientation of the thin film structure is also clear from the analysis of the a> scans curves around the (002) peak for each energy of the ion beam (i.e., rocking curves). The curves, shown in Figure 3.24, reveal a clear evolution from an asymmetrical shape at 100 eV ion energy to a symmetrical shape as the ion energy increases to 300 eV. These two sets of results clearly indicate that the ion energy has a positive effect in inducing the preferential growth of the (002) planes parallel to the surface plane. A direct dependence between the preferential degree of orientation and the ion energy has also been found for metal nitrides. Thus, for example, for titanium nitride thin films prepared by nitrogen ion bombardment of evaporated titanium, it has been found that the (200) texturing degree increases in detriment of the (111) orientation as the energy increases from 0.3 to 1 keV (Ensinger, 1995). This tendency finds a relatively flat maximum for ion energies between 1 and 10 keV. Above this energy, the films start to lose very slowly the preferred (200) orientation. This behaviour suggests that the degree of orientation depends on ion energy but not, however, in a simple way.
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(a)
^15k
145
ff, km beam energy
I
i
3
H
*?10k-
•e s. »
5k-
I 20
-r30
40
50
60
70
80
20 (deg)
Figure 3.24. Effect of ion beam energy on the texture of Ce02 thin films deposited by magnetron IBAD. a) X-ray 0/28 scans; b) CeOa (002) rocking curves. Reproduced from Gnanarajan (1999) with permission.
A common observation made with many compound thin films is that texturing increases up to ion energies around 1000 eV, while above 10 KeV a small decrease in the degree of orientation is generally found. At high ion energies, the structural damage increases while the sputtering yield decreases (cf., sections 3.7 and 1.7.1). These two factors tend to weaken the processes contributing to the preferential orientation and therefore favour certain randomness for atom rearrangements as will be discussed below. This tendency has been commented in the paragraph above when discussing the behaviour of TiN. The ion to atom arrival ratio has also a positive influence in favouring the degree of preferential orientation. Figure 3.25 shows the texture coefficient of the
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(200) plane for TiN grown under N* ion bombardment for two different energies of the ion beam (Ensinger, 1995). At small I/A ratios, the films exhibit a preferred (111) orientation for 6 keV beam energy. However, a complete texturing according to a (200) orientation occurs for an I/A ratio of around 0.5. A similar type of influence of I/A is also observed for ions of 30 keV, although in this case higher I/A ratios are required to induce a significant degree of orientation. Similar tendencies have been reported for other type of thin films such as metals, oxides, etc. In general, when IBAD films are prepared at low ion irradiation intensities, crystallisation of the most energetically stable plane structure is induced. TiN has an fee structure where the (111) planes are the most densely packed and, therefore, have the minimum free energy. Accordingly, (100) texturing only occurs at higher ion doses. Under these conditions, the growth would be under kinetic control, rather than under thermodynamic control, the stabilisation of structures with higher internal energies therefore being possible. 1.0
TiN N+ / */
6keV a
0.6
S 0.4
*
I
/ './
//
/*"""" ,J»
30keV
/ ..*'
0.0 -—1 0.0
/ 1
1
0.2
1
1
0.4
1
1
1
0.6
1
0.8
L.
,
1.0
arrival ratio l/A Figure 3.25. Orientation coefficient (X-ray diffraction peak height ratio) as a function of arrival ratio I/A of TiN deposited under nitrogen ion bombardment with two different ion energies. Reproduced from Ensinger (1995) with permission.
Another common tendency found in IAD thin films is that heavier ions are more effective that light ions in inducing texturing of the films. This experimental finding stresses the importance of the momentum and electronic energy transfer processes for the control of preferential orientation phenomena. A similar conclusion can be drawn from the observation that in series of compounds where the atomic mass of one of the constituent elements is varied, the degree of
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preferential orientation is smaller for the compounds with the heavier elements (Dong et al., 1998; Zhang et al., 1998). Thus, for example, the texture evolution from a (111) to (200) in preferentially oriented TiN thin films occurs more effectively under Ar+ than N* bombardment, as expected from the atomic masses (40 and 14, respectively), of these two ion species (cf., section 1.8.2). Moreover, for equivalent experimental conditions, the texture coefficient of the (200) planes of titanium nitride is 0.8, while it is only -0.3 for tantalum nitride. A less effective momentum transfer in this latter case because of the higher atomic mass of tantalum is likely to be the reason for the smaller texture coefficient found for tantalum nitride thin films prepared by IAD.
3.9.4. Models for texture development Taking into account the aforementioned evidence relating the degree of preferential orientation to the magnitude of the momentum transferred by the impinging ions to the target, initially accepted explanations accounting for texturing effects during thin film growth rely on the idea of the preferential sputtering of certain planes. Within this scheme, only the planes with the smallest sputtering rate would develop in detriment to those with a higher sputtering probability. This concept was initially formulated by Dobrev (1982) and has been extensively used in the literature to account for the preferential growth of crystal planes in many thin film compounds grown with IAD techniques. The basic idea of this model can be schematically explained by considering that when the impinging ions face a family of planes that define well-aligned channel structures, they can penetrate more deeply into the bulk without undergoing collision events. By contrast, if the ion beam faces a family of planes with a closed packed structure where there are no channels parallel to its direction, collisions will already occur with the first plane atoms. Then, preferential erosion of that family of planes would occur, while the planes with channel structures aligned parallel to the beam direction will be preserved. This idea is illustrated in Figure 3.26 showing a scheme of the atoms of the (111) and (002) planes of CeOz projected on the surface normal to the ion beam direction. Just based on geometrical arguments of occupation of surface, it is clear that the probability of collisions of ions impinging perpendicular to the (111) planes will be higher than that on the (002) planes (note that in the cubic Ce0 2 structure the (002) and (200) planes are equivalent), thus favouring the development of a (002) texture in these thin films. According to the model of preferential sputtering, during the initial stages of nucleation and thin film growth, grains with the more favoured plane orientations
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can grow relatively undisturbed either in the direction of the ion beam or laterally. The growth according to this lateral dimension is favoured to the detriment of neighbouring grains initially facing unfavourable crystal planes towards the ion beam. According to Dobrev's original suggestion, in a random or only slightly oriented polycrystalline film there will be some crystallites with the most favoured channel directions coinciding with the ion beam direction. In these crystals the energy lost by collisions will be smaller than in the neighbouring crystallites with other orientations. Upon bombardment, the oriented crystallites will remain cooler and will serve as recrystallisation centres of the adjacent regions in an ion beam activated recrystallisation process.
?o®o®o?o? 5®ofo?ofo ®^®^®^®A® •@ # @ # ©% #
°.0.0.0.0.0 o#o#o#o#o#o o o o o o o o o o o o o
Figure 3.26. Projection of the (111) (left) and (002) (right) planes of Ce0 2 on the surface normal to the ion beam direction, respectively, where the small solid circles represent the Ce and the hollow big circles the O atoms, respectively. The dashed circles represent oxygen atoms which are underneath the position of Ce atoms.
Dobrev's model seems to apply rather well to many compounds. In metals, texture development upon ion bombardment can be predicted by considering the simple criteria of minimum free energy of planes and the easiest channelling directions. Thus, for the different metal structures, preferential texturing in IBAD thin films should be expected to occur according to: fee: bec: hep:
(111) (non-bombarded) -> (110), (100), (111) (110) (non-bombarded) -» (111), (100), (110) (0002) (non-bombarded) -»(1 HO), (0002)
Bradley et al. (1986, 1987) proposed a model to explain the development of preferred orientations in IBAD films that relies on the difference in sputtering yields according to different orientations rather than on a reorientation process during crystallisation.
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Molecular dynamic simulations have been also used to describe preferential growth phenomena in thin films. The calculations confirm the intuitive ideas of the preferential sputtering model outlined above. Figure 3.27 shows the evolution with time of the atomic structure of a bicrystal film grown by IBAD techniques. One of the crystals is oriented with the [111] direction parallel to the z~ axis and the [110] and [112] directions along the x and y axes, respectively. The other crystal is oriented with the [110] direction parallel to the z-axis and the [ 110 ] and [001] directions along the x and y axes (Donget al., 1998). By this calculation the ion beam was oriented perpendicular to the nominal surface. This direction is parallel to the [110] oriented grain, but coincides with the non-channeling [111] direction of another grain. The results of the molecular dynamic simulations depicted in the figure confirm the differences in the growth rate of crystals with channelling and non-channelling directions and the subsequent appearance of shadowing effects of one grain with respect to another. The calculations also reveal that differences in ion damage of the two grains provoke a recrystallisation-like grain boundary migration that favours the growth of the channelling grain in detriment to the non-channeling one. By this calculation, it could also be proven that the sputtering yield from the [110] oriented grains was smaller than that of the [111]
(U)
Figure 3.27. Molecular dynamics simulation of the IB AD growth process of a bicrystal system. The ion beam is oriented normal to the bicrystal surface, such that it Is aligned with the [110] channelling direction of the [110] oriented grain (left) and with the no channelling direction of the [111] oriented grain (right), a) to d) represent structures of the bicrystal after increasing periods of time. Reproduced from Dong et al. (1998) with petmission.
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ones. This effect can be considered as an additional factor contributing to [110] texturing of these films.
3.9.5. Biaxial in-plane orientation To improve the thin film homogeneity, a common practice during thin film deposition is to rotate and/or move laterally the sample holder. In this way, differences in thickness, texture or microstructure arising from lateral inhomogeneities in the profile of incoming particles of the material to be deposited and/or in the beam profile are avoided. When such a technique is used, texture development can only be observed for planes parallel to the surface and no biaxial orientation phenomena can be detected. On the other hand, ion bombardment during IBAD at angles inclined with respect to the surface normal usually produces a change in texturing as compared with the effect of an ion beam perpendicular to the surface. This phenomenon is particularly clear in metals where a progressive evolution of texturing from one family of planes to another can be induced easily by tilting the beam direction in respect to the surface normal. This has been observed, for example, for Cr (10 nm)/CoCrPt (49 nm) films deposited by DIBS. In this system, the intensity ratio between the (1010) and (0002) diffraction peaks changes from about 0.35 to 0.10 when bombardment angles vary from 40° to 90° with regard to the surface (Leng et al., 1999). Thanks to the new developments in the technology of ion and evaporation sources, homogeneous molecule and ion beam profiles can be obtained over large surface areas and rather homogenous films are, therefore, obtained even on immobile substrates. This new technology has opened up the possibility of inducing biaxial thin film orientations by tilting the beam direction with respect to a fixed substrate position. In fact, when the ions impinge perpendicular to the sample surface, only texturing along the film growth direction can occur and, in principle, the grains will have a random azimuthal orientation. By contrast, off-normal bombardment at fixed sample position may render biaxially aligned thin films. In such cases, the grains not only shear one crystallographic axis fixed with respect to the growth direction, but also other axis (and consequently all the axis). This situation is equivalent to that of an oriented single crystal. Biaxial thin film texturing by IBAD films has been observed for many type of materials including nitrides and oxides of high technological interest (Ensinger,
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1994, 1998). An interesting example of biaxial alignment is that of the ytria stabilised zirconia systems (YSZ). Whatever the impinging angle of the ion beam, IB AD YSZ films depict a preferential orientation of the [100] direction perpendicular to the sample surface (i.e., the (100) family of planes are parallel to the surface). This means that the crystallographic c axes of the crystallites are parallel to the surface normal. However, [111] or [110] biaxial alignments can be induced depending on the incident angle between the impinging ions and the substrate. The value of other parameters such as substrate temperature or the ion to molecule arrival ratio (i.e., I/A) may also affect the type of biaxial alignment. Taking XRD polar plots can monitor biaxial orientation of YSZ thin films grown by IBAD. Figure 3.28 shows the (111) x-ray pole figures of (100) biaxially aligned YSZ films with (110) or (111) in-plane orientations with respect to the ion beam. As indicated in the figure, the specimen were oriented so that the projection of the assisting ion beam direction corresponded to (|)=0o. In both cases, the beam was forming an off-normal angle with regard to the substrate. Polar plots in Figure 3.28 are consistent with an in-plane arrangement of crystal axis as schematised in Figure 3.29 for preferential growth of either (111) or (110) planes in respect to the projected beam direction. Note that the (111) planes in the (110) in-plane oriented film do not face the projection of the beam direction, but form a 45° angle with it. This effect is clearly detected in Figure 3.28 by the tilting of 45° found for the (111) pole positions with respect to the projected beam direction.
Figure 3.28. ( I l l ) X-ray pole figures of (100) biaxially aligned YSZ films with (a) (110) in plane orientation and (b) (111) in plane orientation, with respect to the assisting ion beam. The specimen were oriented so that the projection of the assisting ion beam direction corresponds to (|) = 0. Reproduced from Ressler et al. (1997) with permission.
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110)
(ioo)
V
(100)
Figure 3.29. Schematic representation of the orientation of crystallographic axes of IB AD YSZ films (100) biaxially aligned with (111) and (110) in-plane orientations.
Biaxial in-plane alignment of YSZ thin films were first reported by Iijima et al. (1993) by bombardment at an angle of 55° from the substrate normal. This angle corresponds to the direction of a [111] axis in the YSZ unit cell with [100] fixed normal. This finding seems to support the selective sputtering model since the films appears to grow preferentially along the [111] channelling axis. However, later studies have shown that by changing the ion to molecule arrival ratio and the ion bombardment angle it is possible to induce either [111] or [110] biaxial growth. These findings have suggested that other factors besides ion channelling may contribute to biaxial alignment in IBAD films. Anisotropic damage has been proposed as a reason for the preferential development of a certain biaxial alignment (Ressler et al., 1997). Thus, it has been argued that the main factor that favours a given orientation is the relative capacity of the planes (110) or (111) to withstand ion damage. In particular, it was demonstrated that for relatively higher ion bombardment doses the (111) plane is more damage tolerant than the (110) one. The reason for such an enhanced resistance seems to lay in the fact that preferential sputtering of oxygen is greater from the (110) planes than from the (111) ones. In agreement with this idea, (111) biaxial orientation is generally observed for higher beam energies and ion to molecule (I/A) ratios. It is also observed that the change from (110) to (111) in-plane orientation typically occurs at low I/A ratios and low bombardment angles. Under these latter conditions, the energy and momentum are dissipated closer to the surface and it seems that the 0 = ions are preferentially sputtered from the (110) rather than from the (111) planes.
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3.9.6. Applications of textured thin films Growth of textured thin films with preferential plane orientations has found interesting applications in many emerging scientific and technological fields. Many thin films are epitaxially grown on lattice-matched single crystal substrates. Very often, these substrates are disadvantageous because their high dielectric constants, low thermal conductivities or coefficients of thermal expansions do not match those of the films. Additionally, other practical shortcomings are the price and the small size of the single crystal substrates available on the market. The possibility of preparing biaxially aligned IBAD films on polycrystalline or amorphous substrates has opened up an interesting route to growing oriented thin films of technological interest on any kind of substrates. In this case, the biaxially oriented IBAD film acts as an intermediate buffer layer on top of which the thin film grows in register. An interesting example of this type of application is the growth of YBa2Cu307.g (YBCO) superconductors on YSZ buffer layers. The use of these buffer layers has not only proved them to be very effective diffusion barriers between the YBCO and the substrate, but also to induce a well oriented grain texture in the YBCO film that contributes to preserve its superconducting properties. Many other oxide and nitride materials besides YSZ have shown biaxial alignment when deposited as thin films with IBAD procedures. Ce0 2 , Ti0 2 , TiN and other metal nitrides, etc. are some of the materials referred to in the literature as presenting this characteristic. Their use as buffer layer substrates to epitaxially grow other type of thin films is clearly an application that will expand in the future. Preferentially oriented metal thin films have also proved to be of great interest because the unusual properties depicted by such systems. Thus, certain textures in metal thin films have been correlated with the development of magnetic properties of technological interest. Permanent magnet layers of Cr/CoCrPt thin films grown by IBAD can be prepared with either ( 1 0 1 0) or (1 1 2 0) textures, with the largest coercivity values being obtained when a strong (10 10) XRD peak is obtained (Hedge et al., 1999). Other properties such as electrical resistivity seem also to depend on the degree of alignment in textured metal thin films.
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3.10. Influence of ion assistance on thin film stress The magnitude of the stress in a thin film is a critical parameter that determines its possible applications. Thus, for example, excessive stress may lead to delamination of the film from the substrate. Other characteristics such as type of crystallographic structure or mechanical stability can be also influenced by the stress. In this section, we will address some general questions about the stress in thin films and the importance of ion bombardment in modifying the stress state and the advantages in terms of practical applications that can be obtained from this feature.
3.10.1. Basic concepts on stress The concept of stress refers to the tension required to produce a certain deformation or strain in a solid. Stress and strain are related by a relationship of the type: a = Ye
(3.6)
where a is the stress or force applied by unit area and £ is the strain or deformation. Kis the Young's modulus. Units of stress are that of pressure (i.e., GN m"2 or GPa). A material subjected to stress undergoes a change in its dimensions. This change is called strain. Tensile stress increases the length of the material. The stress is then considered to be positive. Compressive stress produces a shortening of the dimension of the solid and is considered to have negative sign. An interesting parameter for the discussion of stress and strain is the Poisson ratio, usually referred to as v. This parameter, a positive number for most materials, relates to the expansion or shrinkage of the width of a solid with the increase/decrease of its length when it is subjected to a tensile/compressive stress. For most materials it holds that:
3.10.2. Distribution of stress between substrate and thin film For a thin film deposited on a planar substrate, the stress experienced by the thin film is biaxial, i.e., it acts along the two principal axes in the plane of the film. In
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155
principle, for such geometry, there is no stress in the direction perpendicular to the film free surface, although there is strain in both normal and in-plane directions. An idealised representation of the stress distribution between thin film and substrate is shown in Figure 3.30 for the final part of the thin film/substrate system.
Neutral "plane
substrate
Figure 3.30. Development of compressive stress in a thin film.
In-plane stress in the film changes sign at the interface with the substrate, then passes through zero in the so-called neutral plane and reverses again at the other side of the substrate. At equilibrium, the moment of the force produced by the stress in the film must be equal to that produced by the stress in the substrate (Tu et al., 1992). Bending of the substrate occurs to get a balance between these two bending moments. Figure 3.30 corresponds to a thin film subjected to compressive stress by the substrate. This case might be that of a thin film with a large thermal expansion prepared at low temperature below ambient conditions and then heated up to room temperature. Provided that good adhesion exists between substrate and thin film, no thermal expansion is possible for the film when the system is put at room temperature and then, the substrate bends under the action of the different stress moments developed in the system. Bending of the substrate/thin film system can be used to determine experimentally the magnitude of the stress developed in the film. For this purpose it is very common to use the Stoney's formula (Stoney, 1909):
a =
(3.8)
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Low ENERGY ION ASSISTED FILM GROWTH
where ts and tf are the substrate and thin film thickness, respectively, and r is the curvature radius of the substrate. This formula relates the stress to the bending degree of the substrate. The curvature of the substrate can be measured by many methods, including laser interference, stylus profiling, capacity measurements or Xray diffraction (XRD). By XRD the stress is determined in terms of the shift in the position of the diffraction peaks. The stress in thin films increases with their thickness. When the accumulated stress reaches a critical value delamination can occur.
3.10.3. Thermal stress in thin films In thin films there can be several sources of stress. One is the thermal stress due to the difference in thermal expansion coefficients of the coating and substrate materials. Films develop thermal stress when they are prepared at high (low) temperatures and then the system is cooled down (heated up) to room temperature. During preparation at a certain temperature substrate and thin film will be at equilibrium so that the system can be without stress. However, by changing the temperature above or below that of deposition, lattice dimensions of substrate and thin film usually evolve differently, thus leading to the development of stress. Thermal stress can be approximated by the expression: ath=Y(af-as)(Ts-Ta)
(3.9)
where Of and as are the average coefficients of thermal expansion of thin film and substrate, Ts is the substrate temperature during deposition and Ta is the temperature during determination of stress (i.e., usually room temperature). For Ts > Ta and Of> ccs< the film is under tensile stress as indicated by a positive value of
3.10.4. Intrinsic stress in PVD thin films Another source of stress in thin films is the so-called intrinsic stress. It accounts for the component of measured stress that cannot be attributed to thermal stress. The internal stress is due to the accumulation of crystallographic flaws, slight
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
157
displacements of the positions of the atoms within the lattice or the accumulation of foreign atoms during deposition. When deposition is carried out at relatively low temperatures intrinsic stress can accumulate and dominate over thermal stress. This is a very common situation for thin films prepared by IAD methods. Values of the order of 0.1 - 3 GPa, depending on the deposition conditions are very common for thin films deposited by these techniques. In PVD thin films, microstructure and stress are closely related. The three most important variables determining the microstructure of PVD films and the value of the intrinsic stress accumulated during thin film growth are the temperature of the substrate during deposition, the residual partial pressure in the chamber and the incorporation of foreign atoms within the thin film network. Intrinsic stress in PVD thin films usually decreases with the deposition temperature. The reason for that is simple: adatom mobility increases with temperature, thus favouring the removal of defects, lattice distortions or crystallographic flaws that are important sources of stress in this type of film. The columnar microstructure of the majority of PVD thin films, often characterised by a large fraction of open volume and void boundaries, is another factor contributing to the intrinsic stress in PVD thin films. Such voided structures are more abundant when the films are prepared at low temperatures. Intrinsic stress produced by accumulation of defects or by the existence of a large volume of voids has a tensile character. Since atom mobility is a particular property of each material, T/Tm (Tm melting temperature) is typically used as a universal parameter that enables the comparison of stress of thin films of different materials. In thin films prepared at relatively low T/Tm values, the concentration of defects can be at orders of magnitude higher than that accumulated in bulk materials during cold working. The stress generated by these defects can be partially released by post-deposition annealing of defects by heating the films at higher T/Tm. A schematic representation of the evolution of the two components of stress in PVD films as a function of T/Tm is shown in Figure 3.31. It is apparent in this figure that while thermal stress grows with T/Tm, the intrinsic stress decreases (in the diagram the positive sign of the intrinsic stress indicates that in this particular case it has a tensile character) (Thornton et al., 1989). Metal thin films deposited by PVD methods are very sensitive to contamination by residual impurity gases (e.g. 0 2 , N2, H 2 0, etc.) present in the preparation chamber (Misra et al., 2000). Incorporation of O or N atoms as
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Low ENERGY ION ASSISTED FILM GROWTH
impurities in the thin film lattice structure is a common phenomenon whose relative importance increases when the residual pressure in the deposition chamber is poorer. The impurity atoms occupy interstitial positions of the metal lattice and produce lattice distortions. As a consequence, the thin films develop an intrinsic component of compressive stress that usually depends on the residual pressure in the deposition chamber. By contrast, an opposite tendency (i.e., tensile stress) may occur when an open microstructure develops as a result of the high partial pressure existing in the chamber during deposition. The two tendencies can counterbalance in real thin films, making it difficult to predict in advance which one will prevail.
w w • • — •
CO
T/T m Figure 3.31. Evolution of the thermal and intrinsic stress contributions with the T/Tm ratio.
3.10.5. The stress in IAD thin films: Dependence on experimental parameters PVD thin films are characterised by a high concentration of voids and empty space, in grain boundaries or between columns that generally lead to tensile stress. Ion bombardment of growing thin films is a very effective means to modify the stress state of the film. For very low energy of the assisting ions a slight decrease of void size and an increase of the packing density is observed, although not completely dense microstructures are still obtained. Under such conditions of void distribution and microstructre, the tensile stress may reach maximum values. If the ion energy increases further, a dense film with reduced or even zero stress can be obtained.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
159
Further bombardment drives the film to compressive stress. The trend of a stress state evolving from tensile to zero, to finally compressive for high ion bombardment has been observed experimentally in many metal films (e.g. Cr, Nb, W, Cu, Pd, Au, etc.) (Smidt, 1990 and references therein). This tendency has also been predicted by molecular dynamics simulation. Thus, Miiller, for 2D structures of Ni deposited on Ni was able to reproduce the general behaviour of a tensile stress that increases with ion energy up to 25 eV (I/A=0.16) and then decreases to near zero at around 100 eV (Muller, 1986, 1987). As an example, Figure 3.32 shows the experimental stress evolution as a function of the negative bias applied to the substrate for Cr films (150 nm) deposited by DC magnetron sputtering on Si (100) for two different pressures in the deposition chamber (Misra, et al., 2001). This figure shows that as the bias decreases, the tensile stress increases to a maximum of 1.5 GPa for zero bias voltage. Increase of the negative value of bias voltage leads to relaxation and a rapid build up of compressive stress until saturation above -300 V.
• i i i i • • i i i i i i i i i i i i i • i i i i i i i i i • i i
—B—6 mTorr —•—5 mTorr
OH
x T^—LJ^~S.
is o
(Z3
4»
-6 o -5bb -4bb -3bb -2bb -ibb'
6
ioo
Substrate Bias (V) Figure 3.32. Stress in Cr films deposited by MS as a function of the bias voltage applied to the substrate. Reproduced from Misra et al., (2001) with permission.
Few measurements of stress have been reported for dielectric films such as oxides (e.g. Si0 2 , Ti0 2 , Ta 2 0 5 , etc.) and nitrides (Schmidt, 1990 and refs. therein). Detailed studies on the development of stress during growth of DLC and c-BN films by IAD methods will be discussed in Chapter 5.
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Low ENERGY ION ASSISTED FILM GROWTH
The stress state of the films has a clear influence on other characteristics such as adhesion or cracking. For Zr deposited on Pyrex in a DIBS system where the growth was assisted with an end-Hall ion source (cf., Chapter. 2) it was shown that depending on the flux and energy of the assisting Ar* ions, the films de¥eloped compressive and/or tensile stresses which clearly influenced the adhesion and stability of the films (Trigo et al, 1994). An estimation of the stress in terms of the X-ray diffraction patterns indicated that the transition from tensile to compressive stress occurs at energy of about 70-80 eV per deposited Zr atom. Above and below this energy, the developed stress caused characteristic cracking and wrinkling of the films as those shown in Fig 3.33. In thisfigure,panel a) shows an image of a sample deposited at 29 eY/atom. It displays an intensive cracking due to the tensile stress. Panel b) shows aflatand smooth surface corresponding to a sample deposited at 74 eV/atom, while panel c) shows a characteristic wrinkle pattern due to the compressive stress developed in thefilmfor higher energies of ion assistance.
Figure 3.33. Optical micrographs of & samples deposited by DIBS under Ar+ assistance with a) non-assisted, b) 29 eV/atom, c) 74 eV/atom aid d) 120 eV/atoms. Reproduced from Trigo et al. (1994) with permission.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
161
Besides beam energy the ion to atom ratio (i.e., I/A) has a clear influence on the stress state of the films. Thus, whereas reactively evaporated TiN films (1.5 1CT2 Pa N2 atmosphere) show a tensile stress value of 0.5 GPa, TiN films prepared by IBAD with ions of 2 keV show a compressive stress which increased linearly with the ion to atom ratio and the square root of the ion energy (Gerlach et al., 1998). This tendency agrees with the predictions of the model developed by Windischmann, (1987, 1992), which will be discussed in section 3.10.7. This model predicts a rapid increase of the tensile stress until a maximum is reached, followed by a relaxation towards compressive values when the normalised momentum increases. At much higher reduced momentum, the compressive stress saturates or decreases in absolute value. Several authors have observed this behaviour experimentally.
3.10.6. Compressive stress in IAD thin films Ion bombardment of growing thin films is a very effective means of modifying the stress state of a thin film. In general, compressive stress arises when the growing film is bombarded with ions of tens or hundreds of electronvolts. The generation of compressive stress has been explained in terms of atom peening processes produced by the ion bombardment (dUeurle, et al., 1989). Peening causes atoms to become very effectively packed into the ion-assisted film, which, in this way, develops a higher density than the equivalent PVD, thin film. In fact, under ion bombardment atoms may be forced into lattice spaces which otherwise would be too small to accommodate them under thermal equilibrium conditions. This extra occupation of the space produces an expansion of the film outwards from the substrate. However, at the substrate interface the film is not free to expand and the entrapment of atoms causes the appearance of macroscopic compressive stress. This effect is added to the aforementioned removal of voids with respect to PVD thin films. The occurrence of peening effects does not always imply that all IAD thin films are under compressive stress. The state of stress of an actual thin film depends on a series of experimental parameters such as T/Tm (i.e., thermal effects), energy of the ions or I/A ratio and on other factors like the incorporation of noble gases into the lattice or the preferential sputtering of impurity atoms of O and N incorporated in metal PVD thin films during their growth. Referring first to the incorporation of noble gas atoms as a source of compressive stress, MDS calculations have shown that Ar atoms incorporated within a metal lattice induce local fields of compressive
162
Low ENERGY ION ASSISTED FILM GROWTH
stress around these atom positions that lead to the development of macroscopic compressive stress in the film (Fang et al., 1993). It is interesting that the development of compressive stress due to noble gas incorporation may be accompanied by a decrease in the density of the thin film, the opposite to that considered by the peening model. The incorporation of Ar or other noble gas atoms in the lattice of the thin film and the consequent increase in compressive stress may contribute to changes in the crystallographic structure and/or in the preferential orientation of certain crystal planes (Nowak et al., 1999). Thus, for example, stabilisation of the tetragonal crystallographic phase of ZrC>2 in thin films of this material with Ar atoms incorporated within its structure has been related to the development of an intense compressive stress because of the incorporation of such atoms (Holgado et al., 2001). In metal thin films with impurity O or N atoms, ion bombardment may lead to the development of tensile rather than compressive stress or, at least to the compensation of this latter (Misra et al., 2000). This is due to the preferential removal by sputtering of the impurity atoms and the release of compressive stress produced by their incorporation during deposition of the films (Cuomo et al., 1982). At the limit, the stress in the film may even become tensile because of this preferential sputtering.
3.10.7. The stress in IAD thin films: Models Many authors have discussed the evolution of the compressive stress as a function of ion energy and I/A ratio for films deposited by IAD methods. Windischmann (Windischmann, 1987, 1992) presented a model based on the knock-on linear cascade theory of Sigmund (Sigmund, 1981) to determine the fractional number of atoms displaced from equilibrium, which is then related to the volumetric distortional strain and stress. The model predicts that the compressive stress accumulated during deposition can be approximated by the expression:
a = cted Q
.V.
E/2
(3.10)
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
163
where 8 is defined in terms of the sublimation energy of the target, U0, atomic number and atomic mass of the projectile and the target ZpJ and MpJ, as
P
5=
„,
rr^r
(3.11)
Uotl + M . / M ^ Z ^ + Z / V and Q is the elastic energy per mole given by the expression: M Q = Y-—t— (l-v)p
(3.12)
where p is the density of the sample. Eqn. (3.10) indicates that the compressive stress in an IAD thin film depends on the material characteristics (through Q and 8) and is directly proportional to the I/A ratio and the square root of the energy. The dependence of a on (I/A)EI/2 is probably the main feature of Windischmann's model indicating that the compressive stress in a IAD thin film is a process driven by the ion momentum transfer. The model applies rather well to a large set of materials, particularly metals, deposited under bombardment with energetic particles during growth. However, although many experimental results can be rationalized by assuming a direct dependence of a on Em, there are also many other examples where a different dependence has been found. Davis (Davis, 1993) modified Windischmann's model to take into account stress relaxation effects and film deposition rates. These effects were not considered by Windischmann, who simply assumed that the stress is caused by the volumetric distortion resulting from atoms that have been displaced from their equilibrium positions after collisions with incident ions or recoiling atoms and remain frozen in the new positions. Davis's model assumes that knocked-on atoms of the film that are implanted below the surface cause the compressive stress. Relaxation is incorporated through the thermal spikes that are induced by the impinging ions. As a first approximation, the strain is assumed proportional to the fraction of implanted atoms in the films. The stress a is given by:
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Low ENERGY ION ASSISTED FILM GROWTH
o(E)>
( Y 1-v
(3.13)
5/
J
A/I + kEA
Where A is the rate per unit area with which atoms are added to the growing film and / the ion beam density; A can be expressed as A—d-p, where d is the growing rate in film thickness and p is its atomic density. K is a constant given by K = 0.016 (p E0'5/3, where E0 is a parameter that can be considered as effective excitation energy and would indicate that the ions must have certain energy above a certain threshold value to become effective in producing atom displacements.
A/l=0.5
A/l=1
3.
A/l=2
I
A/l=10
3
20
40
60
80
100
E(eV) Figure 3.34. Evolution of the stress according to Eqn. 3.13 as a function of the ion energy for different values of A/I.
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
165
Other models have been developed to account for the evolution of stress as a function of process parameters. Ward and Williams (Ward et al., 1999) developed a model consisting of a finite element simulation of the stress. The model allowed the generation of stress profiles through the film thickness and substrate, showing that the overall stress value in the film follows a similar evolution with E as that predicted by the Davis's model. Usually the models concentrate on the tensile or compressive stress separately, so that whereas the tensile stress models are based on the estimation of the interatomic forces acting on the grain boundaries of the columnar structure (Itoh et al., 1991; Misra et al., 2001), the compressive stress is very often considered to be due to the atomic peening process mentioned above and sometimes as the consequence of gas incorporation in the film. Knuyt et al., (2000) have presented a model which describes the overall behaviour of the residual stress in a film deposited by IAD methods, including the transition from tensile to compressive stress.
3.11. Improvement of adhesion in IAD thin films The first requirement of a coating is that it adheres well to the substrate. Therefore, adhesion is a critical factor in the manufacture of coatings aimed at a wide range of applications. The interest in adhesion between dissimilar materials is increasing following the development of new capabilities of modifying interfaces at nanoscopic and even atomic scales. According to Baglin (1994), a stable bonded interface between a coating and its substrate must, in general, involve an intermediate layer, which should be chemically stable and intrinsically linked to both the film and the substrate. The role played by such an interface layer is to lower the interface free energy in order to maximize the energy of adhesion Waj given by: " ad
=
7film + Ysubst ~ 7m
(3-14)
where y is the corresponding surface free energy for the film, substrate and interface. Although the formation of an interface layer will depend on the system under consideration, in many cases the ion beam processing can induce it. Obviously, the quality of adhesion is specific to the application and includes the requirement that it will remain stable with the passage of time.
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Low ENERGY ION ASSISTED FILM GROWTH
The promotion of the adhesion by post-deposition bombardment and ion mixing using high-energy ions has been extensively studied. An excellent review of the subject and a list of references to experimental data is given by Baglin, (1987), and the reader is referred to it for more details, because it is beyond the objectives of this book. Here, we will focus the discussion on the use of low energy ion beams and IAD methods. The approach of using low energy ion bombardment to influence the adhesion of a coating with the substrate is based on their capabilities to control and modify the intrinsic stress of the coating (cf., section 3.10), the interface morphology and the interface chemistry, or to induce the removal of surface contaminants by sputtering, (Baglin, 1994). Sputtering of the substrate, previous to the deposition, provides a well controlled means of modifying the substrate surface to produce a series of effects, e.g. sputter-cleaning of contaminants, modification and new formation of surface bonds, the increase of the roughness and effective surface area, changes of the surface composition and surface chemistry when using reactive gases, etc. If the ion bombardment continues during the growth (IAD deposition methods), the growing films become denser and even stress free and, therefore, have improved adhesion. The beneficial influence of sputter cleaning before deposition to eliminate contaminants from the substrate surface is well established. The results reported by Cailler et al. (1993) constitute a good demonstration of that influence. The data show the successful adhesion of a 200 nm thick Cu-film on a polished carbon steel after removal of the surface oxide layer with 600 eV Ar+-ions. The film adhesion was observed to increase by a factor of 20, as compared with the untreated surface. In the case of a surface compound, sputtering not only removes contaminants, but also induces significant chemical and compositional modifications, so that adhesion can be also modified. This kind of effect has been reported for the Cu-Al 2 0 3 system. A study of the influence of in-situ bombardment of sapphire with 500 eV Ar+, previous to the evaporation of Cu showed that the peel strength of the coating depended on the ion dose necessary to form some mixed AlO-Cu complexes at the interface, which clearly improved the adhesion of the coating (Baglin et al., 1987). The formation of such interface layer was evidenced by XPS and AES. Improvement of adhesion attributed to some kind of chemically induced effects, has also been reported for IBAD gold films deposited on glass by
EFFECTS INDUCED BY THE ION ASSISTANCE OF FILM GROWTH
167
using oxygen or oxygen + argon ion beams. The reported enhancement of adhesion induced by the oxygen bombardment was up to 200-400 times more than nonassisted or only Ar-assisted coatings (Martin et al., 1985). The metallisation of polymers is a research area of great interest, where the adhesion constitutes a problem that can be afforded by IAD methods (cf., section 4.5). The effect of sputtering cleaning on adhesion and chemical bonding at the metal/polymer interface has been recently studied by Fujinami et al. (1998) for Ti deposited on PE and PTFE after sputtering cleaning. Adhesion of Ti films to these two polymers was improved by sputtering cleaning, but the effects on chemical bonding at the interface depended on the chemical nature of the substrate. It was observed that while for Ti/PE, the pull strength of the Ti film increased with the ratio of C-Ti bonds (as determined by XPS) at the interface, for Ti/PTFE, where both the C-Ti and F-Ti bonds are observed to form by XPS, there was no correlation between the formation of such bonds and the adhesion enhancement. In the latter case, the adhesion improvement was attributed to the roughening of the PTFE surface by Ar-bombardment, as corroborated by SEM images. Figure 3.35 shows the pull strength of Ti films on the two polymers, as a function of Ar+ dose. In both cases the substrates were first sputter-cleaned before the metal was deposited. Although both improved the adhesion, the effectiveness was higher for PE than for PTFE.
i 0.0
04
0.8
1.2
Ion Beam Current [A|
Figure 3.35. Bond strength for a Ti film deposited on PE ( • ) and PTFE (O) as a function of the ion beam current. Reproduced from Fujinami, et al. (1998) with permission.
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Low ENERGY ION ASSISTED FILM GROWTH
In summary, improvements of thin film adhesion by low energy ion bombardment seem to arise from three effects, sputter cleaning of the surface, ion induced mixing and chemical modification of the interface and sputter induced roughness.
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De Hosson, J.Th. M., Kooi, B.J., Microstructure and Properties of Interfaces between Dissimilar Materials; Nalwa, H.S., ed., Handbook of Surfaces and Interfaces of Materials, vol.1, p. 2, Academic Press 2001. DUeurle, F.M., Harper, J.M.E., Thin Solid Films 171 (1989) 81. Dong, L., Srolovitz, D.J., J. Appl. Phys. 84 (1998) 5261. Dobrev, D., Thin Solid Films 92 (1982) 41. Durand, H.A., Sekine, K., Etoh, K., Ito, K., Kataoka, I., Thin Sol. Films 336 (1998) 42. Durand H.A., Sekine, K., Etoh, K., Ito, K., Kataoka, I., /. Appl. Phys. 84 (1998) 2591. Eckstein, W., Biersack, J.P., Z. Phys. B63 (1986) 471. Ensinger, W., Surf. Coat. Technol. 65 (1994) 90. Ensinger, W., Nucl. Instru. Meth. Phys. Res. B 106 (1995) 142. Ensinger, W., Kiuchi, M., Surf. Coat. Technol. 94/95 (1997) 433. Ensinger, W., Surf. Coat. Technol. 99 (1998) 1. Family, F., Vicsek, T., J. Phys. A L75 (1985) 18. Fang, C.C., Jones, F., Prasad, U., J. Appl. Phys. 74 (1993) 4472. Fujinami, Y., Hayashi, H., Ebe, A., Imai, O., Ogata, K., Mater. Chem. Phys. 54 (1998) 102. Gerlach, J.W., Kraus, T., Sienz, S., Moske, M., Zeitler, M., Rauschenbach, B., Surf. Coat. Technol. 103/104(1998)281. Gerlach, J.W., Schwertberger, R., Schrupp, D., Rauschenbach, B., Neumann, H., Zeuner, M., Surf. Coat. Technol. 128/129 (2000) 286. Gnanarajan, S., Savvides, N., Thin Solid Films 350 (1999) 124-129. Grigorov, G.I., Martev, I.N., Langeron, J.-P., Vigues, J.L., Thin Solid Films 161 (1988)249. Guarnieri, C.R., Offsey, S.D, Cuomo, J.J., in Handbook of Ion Beam processing Technology, eds. Cuomo, J.J., Rossnagel, S.M., Kaufman, H.R., Noyes, Park Ridge, NY1989,p.l89. Hedge, H., Wang, J., Devasahayam, A.J., Kanarov, V., Hayes, A., Yevtukhov, R., Bozeman, S., Anderson, P., Tabat, N., Ryan, P., J. Vac. Sci. Technol. B 17 (1999) 2186.
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Hirvonen, J. K., Mater. Sci. Report 6 (1991) 215. Holgado, J.P., Espinos, J.P., Yubero, F., Justo, A., Ocana, M., Benftez, J., GonzalezElipe, A.R., Thin Solid Films 389 (2001) 34. Hultman, L., Helmersson, U., Barnett, S.A., Sundgren, J.-E., Greene, J.E., J. Appl. Phys. 61 (1987) 552. Iijima, Y., Onabe, K., Futaki, N., Sadakata, N., Kohno, O., J. Appl. Phys. 74 (1993) 1905. Itoh, M., Hori, M., Nadahara, S., J. Vac. Sci. Technol. B 9 (1991) 149. Karakaraju, S., Mohan, S., Sood, A.K., Thin Sol. Films 305 (1997) 191. Kim, K.-S., Shim, H.-S., Kim, S.-H., J. Crystal Growth 212 (2000) 74. Kiuchi, M., Nucl. Instr. Meth. Phys. Res. B 80/81 (1993) 1343. Koch, Th., Ziemann, P., Thin Solid. Films 303 (1997) 122. Knuyt, G., Lauwerens, W., Stals, L.M., Thin Solid. Films 370 (2000) 232. Kuratani, N., Murakami, Y., Imai, O., Ebe, A., Nishiyama, S., Ogata, K., J. Vac. Sci. Technol. A 15 (1997) 3086. Lee, Ch-Ch., Hsu, J.-Ch., Wei, D.T., Lin, J.-H., Thin Sol. Films 308/309 (1997) 74. Leng, Q., Mao, M., Miloslavsky, L., Simion, B., Hung, C-Y., Qian, C , Miller, M., Basi, R., Tong, H.C., Wang, J., Hegde, H., J. Appl. Phys. 85 (1999) 5843. Lifshitz, Y., Lempert, G.D., Grossman, E., Phys. Rev. Lett. 72 (1994) 2753. Ma, Z.Q., Kido, Y., Thin Sol. Films 359 (2000) 288. Martin, P.J., Sainty, W.G., Netterfield, R.P., Vacuum 35 (1985) 621. Mayer, J.W., Tsaur, B.Y., Lau, S.S., Hung, L.S., Nucl. Instr. Meth. 182/183 (1981)1. Misra, A., Nastasi, M., J. Vac. Sci. Technol. A 18 (2000) 2517. Misra, A., Nastasi, M., Nucl. Instr. Meth. Phys. Res. B 175/177 (2001) 688. Miiller, K.H., J. Appl. Phys. 59 (1986) 2803. Miiller, K.H., Phys. Rev. B 35 (1987) 7906. Netterfield, R.P., Miiller, H.H., Mckenzie, D.R., Goonan, M.J., Martin, P.J., J. Appl. Phys. 63 (1988) 760. Nevot, L., Croce, P., Revue Phys. Appl. 15 (1980) 761. Nordlund, K., Ghaly. M., Averback, R.S., J. Appl. Phys. 83 (1998) 238.
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Nowak, R., Yoshida, F., Margiel, J., Major, B., J. Appl. Phys. 85 (1999) 841. Paine, B.M., Liu, B.X., "Ion Beam Mixing"; Itoh, T., ed. "Ion Beam Assisted Film Growth", Elsevier, Amsterdam, 1989, p. 153. Pichon, L., Girardeau, T., Straboni, A., Lignon, F., Guerin, P., Perriere, J., Appl. Surf.Sci. 150(1999)115. Paturand, C , Forges, G., Sainte Chaterine, M.C., Machet, J., Thin Sol. Films 347 (1999) 46. Pranevicius, L., Thin Sol. Films 63 (1979) 77. Reinke, S., Kulisch, W., Surf. Coat. Technol. 34/86; 97 (1997) 23. Sanz, J.M., Soriano, L., Prieto, P., Tyuliev, G., Morant, C , Elizalde, E., Thin Sol. Films 332 (1998) 209. Schneider, J.M., Sproul, W.D., Voevodin, A.A., Mathews, A., J. Vac. Sci. Technol. A 15 (1997) 1084. Sigmund, P., Sputtering by ion bombardment: Theoretical concepts in Sputtering by ion bombardment I: Physical sputtering of single-element solids, Behrisch, R., ed. Topics in Applied Physics vol. 47. Springer-Verlag, Berlin, 1981. Smidt, F.A., Internt. Mater. Rev. 35 (1990) 61. Smith, R.W., Srolovitz, D. J., J. Appl. Phys. 79 (1996) 1448. Sohn, M.H., Kim, S.I., /. Vac. Sci. Technol. A 18 (2000) 1983. Stoney, G.G., Proc. R. Soc. London 82 (1909) 172. Telling, N.D., Crapper, M-D., Lovett, D.R., Guifoyle, S.J., Tang, C.C., Petty, M., Thin Solid Films 317 (1998) 278-281. Thornton, J.A., Ann. Rev. Mater. Sci. 7 (1977) 239. Thornton, J.A., Hoffman, D.W., Thin Sol. Films 17 (1989) 5. Trigo, J.F., Elizalde, E., Quiros, C , Sanz, J.M., Vacuum 45 (1994) 1039. Tu, K.N., Rosenberg, R., eds. Analytical Techniques for Thin films. Treatise on Materials Science and Technology. Vol. 27. Academic Press, San diego 1988. Tu, K-N., Mayer, J.W., Feldman, L.C., "Electronic Thin Film Science for Electrical Engineeres and Materials Scientists", Macmillan Publ. Co, New York 1992, p. 84. Vechten, van D., Hubler, G.K., Donovan, E.P., Correll, F.D., J. Vac. Sci. Technol. A 8 (1990) 821. Volz, K., Kiuchi, M., Ensinger, W., Surf. Coat. Technol. 108/109 (1998) 303.
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Wagner, T.A., Oberbeck, L., Bergmann, R.B., personal communication. EMRS meeting, Strasbourg, June 2001. Ward, D.J., Williams, A.F., Thin Sol. Films 355/356 (1999) 311. Watanabe, Y., Uchiyama, Sh., Nakamura, Y., Li, Ch., Sekino, T., Niihara, K., /. Vac. Sci. Technol. 17 (1999) 603. Windischmann, H., J. Appl. Phys. 62 (1987) 1800. Windischmann, H., Crit. Rev. Solid State Mater. Sci. 17 (1992) 547. Winterbon, K.B., "Ion Implantation Range and Energy Deposition Distributions", 2 Plenum, New York 1975, p. 161. Zhang, F., Zheng, Z., Chen, Y., Liu, D., Liu, X., J. Appl. Phys. 83 (1998) 4101. Zhou, X.W., Wadley, H.N.G., J. Appl. Phys. 87 (2000) 2273. Zhou, X-W., Wadley, H.N.G., J. Appl. Phys. 87 (2000) 8487.
CHAPTER4 APPLICATIONS OF IAD PROCESSING
Ion bombardment during growth influences many microscopic properties of the deposited materials, e.g., film composition, structure, density, grain size, crystallographic orientation, morphology, topography and many other features which have been discussed in Chapter 3. As a consequence of those modifications, the IAD methods have shown their ability to improve a large variety of film properties such as the refractive index, optical absorption, stress, hardness, porosity, magnetisation etc., that depend on the modified structure and microstructure of the films. These modifications make possible to improve the functionality of the growing thin film to obtain better wear- and corrosion-resistant components or coatings with special electro-magnetic or optical properties. Although many of the applications of these modified films are probably protected trade secrets, some examples that have been reported in the open literature are presented in the following. The approach is to show the wide range of applications of IAD processes with the help of some selected examples, where the experimental observations have been correlated with deposition parameters and microscopic effects. Very often, the literature only reports on the improvement of a property or functionality under certain experimental conditions, wrongly given the impression that it is the result of trial and error procedure. In some other cases, the results published in the open literature require a deep critical observation and sometimes an independent confirmation. In the following we will try to avoid this type of report. While the more conventional and extended coating methods, such us electrolytic and electro-less procedures require only modest investments in equipment, but an extensive and intensive process control and empirical experience, the introduction of PVD and IA-PVD based methods is still handicapped by a high investment in equipment and technology. Also the requirement of vacuum makes these methods expensive and more complex. However the possibility of obtaining better properties, even at low temperatures is obviously an advantage that in many cases constitutes a requirement of the application. The use of IAD methods is still limited to specific areas where there are either quality reasons or a significant added value in the product. Nevertheless, several IA-PVD (e.g., IBAD, PHI, etc.) are 173
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finding an increasing amount of interest in areas where wear and corrosion resistance, optical stability, texture, adhesion, synthesis of new materials, etc., are critical issues (Ensinger, 1998; Celis et al., 1999; Kelly and Arnell, 2000; Mattox, 2000).
4.1. Tribological coatings Modern technology has significantly increased the mechanical thermal and chemical load on the materials used in many technological applications. The coatings are a suitable approach to overcome such overloads and to maintain the performance of the materials. There is a large demand to develop reliable coating systems and technologies. Whereas the coating materials are usually oxides, carbides, nitrides, borides, aluminides and silicides, a great variety of techniques and strategies, including IAD processes, have been developed to modify and control the microstructure of the coating system to improve the functionality of the coating and the performance of the materials. Due to the difficulty of proposing a universal coating and technology for the large variety of tribological applications, the experimental strategies depend very strongly on the operational parameters, so that only final enhancements are actually pursued (e.g., adherence between the coating and the substrate, better corrosion protection, higher sliding wear resistance, lower friction, etc.). In fact, a proper selection of a coating-substrate system requires the consideration of many crucial parameters including the coating method, thickness, composition, hardness, adhesion, friction and wear properties, residual stress, thermal expansion coefficients, etc. IAD methods are well known to produce high-density coatings with a noncolumnar structure and good morphology. Therefore, they are techniques that have received a lot of attention in the past decade for coating production. Although the largest interest appears focused on TiN and DLC coatings, there have also been many studies on other materials, e.g., CrN, TiCN, MoS2, BN, etc. In the following we present a series of examples of different coatings, where the use of IAD methods for their deposition has been demonstrated to improve their performance in different applications. Most of the examples are academic demonstrations, whose potential commercialisation is expected to follow very slowly and obviously after the appropriate industrial scaling.
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A primary limitation of PVD methods in general and many of the IAD methods in particular is that they are line-of-sight methods that makes processing of complex shapes difficult, although special designs to overcome these limitations have been suggested in the literature (cf., section 2.3). In fact, PHI has become a good alternative as it enables the application of ion implantation and surface modification at an acceptable cost to complex shapes (Ensinger, 1998).
4.1.1. Hard and wear resistant coatings While CVD is the most widely used technique to produce coated inserts, PVD coatings are gaining acceptance in certain applications such as milling, drilling, threading, etc., which require sharp edges or specific finishing. Industrial coaters commonly use sputtering, e-beam evaporation, and arc evaporation, but in many cases ion assistance is incorporated in the process to profit from the unique advantages of the IAD methods (e.g., low deposition temperature and ability to control the microstructure and to obtain smooth and stress free coatings) (Su et al., 1997; Prengel et al., 2001). Coatings deposited by IAD methods are especially appropriate on substrates such as high temperature alloys and austenitic stainless steels where low deposition temperatures are necessary. The most widely used wear resistant coatings are A1203, TiC, TIN, TiCN and TiAIN, depending on the specific application. TiN, produced either by PVD or CVD deposition methods is certainly one of the most widely investigated coatings. It has found a wide range of commercial and industrial applications, including its use as wear resistant coatings on tool steels or carbides. For these applications, a coating with an optimised structure which yields the highest hardness is required. The possibility of influencing the structure of the coatings by modifying the deposition parameters, e.g., ion energy, I/A ratio and angle of incidence, is what the ion-assisted methods provide. However, an overall improvement and optimisation of the coating properties to obtain a universal coating is not possible, and in general, it is necessary to meet a compromise for each specific application. The result is, rather commonly, that highly crystalline coatings obtained at low I/A show high hardness and good tribological properties, but poor corrosion protection. This property is improved by using more amorphous films deposited at high I/A ratios, although their hardness decreases significantly (Vera and Wolf, 1999).
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Hardness and stress are also two properties mat usually appear closely correlated in coatings deposited by PVD and IAD methods (cf., section 3.10). In fact, that correlation makes difficult the growth of PVD hard coatings with thickness above 6-7 urn, without adhesion problems. As an example, Fig. 4.1 shows the relationship between the micro hardness of 1 urn thick TiN films deposited on silicon by ion assisted arc deposition (IAAD) and the compressive stress as reported by Martin et al. (1999). The stress and hardness were controlled by the [N*]/[Ti] arrival ratio (i.e., I/A) and the energy of the assisting N2+ ions. Figure 4.1 shows that the use of the hardest coating will be at the expense of introducing high compressive stresses and therefore a poorer adhesion and a thickness limitation, so that the coating may become ineffective for many wear applications.
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It is well known that stainless steel has a very good corrosion resistance, but its wear resistance is relatively poor due to its low hardness. However, it has been found that nitridation of steels by PHI hardens and improves the wear resistance of its surface without losing their high corrosion resistance. The nitrided steel becomes then well suited for tools if the process temperature is controlled appropriately
APPLICATIONS OF IAD PROCESSING
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(Ensinger 1998). Up to 450°C the corrosion resistance improves significantly due to the formation of a nitrogen expanded austenite phase. However if the PHI method is performed above 500°C chromium nitride precipitates and the expanded austenite phase transforms to martensite, leading to a dramatic reduction in corrosion resistance (Ensinger, 1998). The development and improvement of properties and performance of coatings for tribological applications is occurring in multiple directions. Multicomponents and graded coatings based on the addition of light elements (e.g., B, Al, Si, etc.) to TiN and TiC, e.g., (Ti, Al) N and Ti (B, N), have attracted an increasing interest as a means of obtaining wear protective coatings with higher oxidation resistance at elevated temperatures and an improved performance in machining operations. These coatings produced by cathodic arc and IBAD, are real alternatives to TiN. However, it turns out that the wear behaviour of these coatings strongly depends on the composition and the degree of improvement, as compared with TiN, is highly dependent on the working parameters of the tool. Multi-layers, including soft and lubricating materials, nano-structured multi-layers (i.e., thickness of the order of nm) and nano-crystalline coatings are also being considered as promising developments (Jehn, 2000). Obviously IAD methods are employed in most of those developments, because of the capabilities offered by them of controlling the microstructure of the coating and of performing depositions at low temperatures. Super-hard and low friction DLC coatings deposited at low temperatures are currently of great interest for wear protection and friction reduction. However, their intrinsic stress and poor adhesion limit the coating thickness and therefore, their potential applications (cf., section 5.6). These handicaps are specially pronounced when deposited on soft substrates such as steels.
4.1.2. Solid lubricant coatings Liquid lubricants have been used for centuries to facilitate sliding in machining. However, there is currently an increasing demand for dry machining, just because the presence of liquids is not recommended or it is even forbidden because of environmental reasons. Accordingly, new coatings and deposition methods as well as post-deposition treatments have been developed in recent years. Depending on applications, friction coefficients ranging from 0.4 to 0.01 are commonly desired.
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Among the most used solid lubricants we find soft metals (e.g., In, Ag, Pb and Au) and lamellar solids of the type MoS2 and WS 2 (Hirvonen et al., 1996). From them, MoS2 is probably the coating that has received the highest attention for vacuum applications. Although usually deposited by magnetron sputtering, ion beam sputtering and IAD methods are also being employed. Due to the preferential sputtering of sulphur some care must be taken to achieve the deposition of stoichiometric M0S2 coatings. Recently, two new coatings have been developed at TEER Coatings using IAD methods, MOST® and Graphit-iC® (Fox et al., 2000, Renevier et al., 2000). The first is a MoS2/Ti composite produced by unbalanced magnetron sputtering or IBAD. It is harder and much more wear resistant and less sensitive to humidity than MoS2. Due to a significant ion bombardment during deposition the structure is amorphous or constituted by very small crystallites. Graphit-iC has been shown to consist of CrC-C multi-layers where the hardness is provided by the ion-bombarded carbon. Apparently, the effect of the Cr incorporation is to reduce the brittleness. In spite of the ion bombardment the bonding is mostly sp2. The wear is poor in dry nitrogen but a small amount of water vapour leads to good wear properties (Renevier et al., 2000). For applications at high temperature, i.e., above 1000°C, there are rather few available coatings. A solid solution of CaF2-BaF2 seems to be a potential candidate. Bhattacharya et al. (1992) have reported the behaviour of IBAD coatings formed by CaF2 and BaF2-CaF2 solid solutions with and without the incorporation of Ag. The best coatings were those with very small-grains obtained with significant ion assistance. Low friction coefficients around 0.3 were determined up to 800°C. Ceramics (e.g., A1203) with a high temperature stability and a high oxidation resistance, constitute also a potential alternative. However, their wear and friction properties hinder their use with common lubricants. Actually, Ag has been suggested as an additive to reduce the friction coefficient and to improve the wear properties of ceramics, even though its adhesion to ceramic substrates is rather poor. Nevertheless, the application of IBAD methods to obtain Ag covered ceramic coatings has shown the possibility of obtaining a significant reduction of wear and good thermal conductivity.
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4.2. Corrosion resistant coatings A common way of preventing corrosion is to coat the material with the appropriate coating. In order to provide adequate corrosion protection, the coating must be uniform, well adhered, pore free and self-healing when physical damage can occur. Electrodeposition of hard chromium on top of a nickel layer is the usual solution to avoid the wear corrosion of steels, stainless steels, copper alloys, etc. exposed to corrosive environments. Due to the presence of cracks these coatings also require some kind of sealing to hinder the corrosion of the substrate. Unfortunately alternative PVD coatings (e.g., CrN, TiN), even with the benefits of avoiding hazardous electrolytes containing Cr, also exhibit micro-pores, which rapidly lead to the corrosion of the substrate. Porosity in electrodeposited and vapour deposited films is, therefore, a well-known failure mode that could be overcome by the more dense and compact IAD films. Improvements in the corrosion and oxidation protection, which are expected from the use of IAD coatings, will be a consequence of their denser microstructure, a better adhesion to the substrate and the capability of reducing their crystallinity and to produce amorphous films. However, up to now only a scarce number of PVD systems for the deposition of metals and oxides on steels are used in practice for commercial applications. Some of them already involve ion beams but only for pre-cleaning purposes. Several groups (Wolf, 1992; Ensinger et al., 1993; Ensinger, 1996; Stippich et al., 1998) have studied IAD coatings for corrosion protection as they can offer low porosity and good adhesion properties as compared with coatings obtained by more conventional methods. Ensinger et al. (1993) have published a comparison of IBAD coatings for wear and corrosion protection with other PVD (sputtering and ion plating) coatings. The concluding remarks of this study pointed out rather clearly that coatings deposited by IBAD at medium energies show better corrosion protection than the coatings deposited by other non-assisted PVD methods. The improvement was especially remarkable when low deposition temperatures were required. The main arguments are a more compact structure and better adhesion of the IBAD coatings. IBAD is presently being commercially used for the production of TiN coatings to protect wet/dry shaver heads and razor blades, from tribological and
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corrosive attack. Miyano and Kitamura (1994) from Matsushita Electric Works developed a process line for TIN coated razor blades of electrical shavers. The substrate was AISI410 steel, which is first sputter-cleaned and then coated by IB AD with a 0.1 urn thick TiN film. The resulting TiN coated shaver blades demonstrated a better corrosion resistance against sweat. They can even be used for wet shaving and directly cleaned in water. Kiyama et al. (1993) from Sanyo Electric Corp. have reported similar results for ZrN coatings deposited by IBAD on electroformed nickel. Gillette ® has also commercialised a razor blade with a thin a-C: H coating deposited by IAD, which is hard and wear resistant and has a low friction coefficient (cf., section 5.6.1). However, it seems that the largest industrial application of these a-C:H coatings is as wear and corrosion resistant coatings in magnetic devices (e.g., disks, tapes, read/write heads) (Bhushan, 1999). The synthesis, characterisation and applications of DLC films will be discussed in detail in Chapter 5. The success of corrosion resistant TiN coatings deposited by IBAD is mainly due to a reduction in film porosity, as compared with films deposited by PVD or even plasma-based PVD techniques that tend to be columnar and textured. The columnar structure and porosity does not necessarily affect the tribological properties negatively, but is clearly detrimental for corrosion protection. Therefore, reduction of the porosity is a requirement in corrosion resistant coatings. Interestingly, several experimental studies (Ensinger, 1998) have evidenced that this protective effect is significantly enhanced by the use of ions at oblique incidence instead of perpendicular to the film surface. The influence of the ion incidence angle on the texture and microstructure of IBAD TiN coatings has been discussed in section 3.9 and studied in detail by several authors (Ensinger 1996; Alberts et al., 1996; Ensinger, 1998). The behaviour against corrosion was measured by Ensinger (1996) in terms of the critical current densities of the iron dissolution from the substrate through the pores of the coating, as measured by cyclic polarization of the samples in a buffered acetic acid (pH 5.6). The results show that the values of these critical current densities relative to the value measured for the coating assisted at normal incidence (0°) are significantly reduced by the use of Ar+ ions at oblique incidence. Furthermore, considering that the relative values of the dissolution currents are also a measure of the porosity of the coating, the results also show that the lowest porosity is obtained for an incident angle of 40°. This value is in accordance with the conditions observed experimentally for the suppression of the well known (100) columnar texture of TiN films grown at normal ion incidence and a consequent densification of the coating. In fact, Miyano et al. (1994) used 2 keV
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off normal nitrogen ions to assist the deposition of TIN films on blades of electrical razors as mentioned above. IBAD techniques with a variable angle of incidence of the assisting ions seem to be a promising deposition method for corrosion protective coatings. Obviously, an optimisation of the corrosion protection requires an appropriate selection of all the deposition parameters, including I/A and ion energy and angle of ion incidence (Ensinger, 1998 a, Vera and Wolf, 1999). In many cases a good corrosion protection has to be coupled with a good wear performance and low friction coefficient. Unfortunately, very commonly, the conditions that optimise one of the properties do not coincide with those that permit the best of the other, so that a compromise has to be met depending on the application. Thus, it is a rather general result that whereas highly crystalline coatings result in higher hardness and lower friction coefficients than the partially amorphous coatings, the amorphous films give better corrosion protection than columnar crystalline films (Vera and Wolf, 1999).
4.2.1. Metal coatings Ensinger (1996) has reviewed the properties of metal coatings for corrosion protection and the reader is referred to that work for details and specific references. In addition to noble metals (Pt, Au, etc.), which are inert in most environments, other metals (e.g., Al, Cr, Ti) are also used as corrosion protective coatings because of their ability to form a compact natural oxide film which is stable and inert in different aggressive media. This enables them to be deposited without pores and with high adhesion. IAD metallic coatings, such as Al, Cr, Ti, Ta, Nb, alloys like Nb-Cr, or semiconductors like Si and Ge have been tested for corrosion protection of other metals, steels and alloys in different media. The general conclusions are rather similar to those mentioned above for TiN, so that the observation of some improvement on the corrosion protection is explained in terms of better adhesion and a denser structure of the coating (e.g., low porosity). However, it seems that although the IAD methods are able to reduce significantly the number of pores in the coatings, the presence of just a few of them enable the corrosion process to proceed further, and although the corrosion rate is reduced, it is not completely eliminated.
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The protection of Al-alloys against corrosion can be achieved by the use of pure Al as coating. The sensitivity of these alloys to temperatures above 150-200 °C makes the use of IAD deposition methods highly appropriate. Ensinger et al. (1993) compared the corrosion protection of an Al-alloy by 2 um IBAD and 3 um ion beam sputtered Al films. The corrosion behaviour was studied according to the salt spray test. This test showed that after 500 hours exposure the uncoated alloy showed severe pitting corrosion with pits of up to several tens of um, whereas the coated alloy reduced the corrosion attack significantly. Figure 4.2 shows the average number of pits per unit area and the pit depth of the uncoated aluminium alloy as compared with samples, which were coated with 3 um of sputtered or 2 um of IBAD aluminium. In spite of its lower thickness, the IBAD coating showed the best results after the test, i.e., the shallower pits as well as the lowest number of them.
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Figure 4.2. Average number of pits and pit depth of an uncoated aluminium alloy after a salt spray test as compared with samples coated with 2 um and 3|im of Al deposited by IBAD or sputtering. Reproduced from Ensinger et al. (1993) with permission.
4.2.2. Oxide and nitride coatings Coatings consisting of oxides (e.g., A1203, Zr0 2 , Cr 2 0 3 , etc.) and nitrides (TiN, ZrN, Si3N4, etc.) have also been proposed for corrosion protection of metals, steels, and alloys. The review by Ensinger (1996) summarises the behaviour of these films as corrosion resistant coatings. As mentioned above, some nitrides deposited by IBAD methods such as TiN and ZrN have found industrial applications.
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Wolf (1992) has reported the behaviour of Al and A1203 coatings deposited by IB AD on 1.7734 steel. A comparison of the corrosion behaviour between the uncoated and 1 |im Al and 1 um A1203 coated steels in 0.1 m NaCl showed that the coated steels reduced significantly the corrosion process with respect to the uncoated steel. Moreover, A1203 showed the more positive potential for pitting corrosion and the process proceeded up to two orders of magnitude slower than for pure Al. In general, it is found that oxide and nitride coatings obtained by IBAD show an excellent corrosion protection of metals, steels and Al-alloys due to their good adhesion and the relatively low number of defects and flaws present in the coating, depending on the thickness and the process parameters such us ion energy, ion angle of incidence and I/A ratio.
4.2.3. Corrosion protection of Magnesium alloys An important area of application of corrosion resistant coatings is the protection of magnesium alloys. Magnesium and its alloys present as advantages their low price, high strength-to-weight ratio and good recycling properties. Magnesium and its alloys have excellent physical and mechanical properties for automotive and aerospace applications (Gray and Luan, 2001), but unfortunately their susceptibility to galvanic corrosion in salt-spray conditions has hindered their use for many of these applications. One method of preventing corrosion is to cover the material with the appropriate coating. There are a number of technologies available for coatings on magnesium and alloys, including electrochemical plating and anodising, conversion coatings and obviously PVD and CVD methods in both the assisted and non-assisted versions. An important advantage of IA-PVD methods is that the deposition temperature can be maintained below 180°C, the stability temperature of many magnesium alloys. Anodisation is one of the conventional methods to produce a mixed MgOMg (OH)2 coating that requires some sealing or painting to reduce the porosity and to improve the corrosion protection. Therefore dense MgO coatings deposited by IBAD could be good candidates for corrosion protection of Mg-alloys. Stippich et al. (1998) have tested 1 urn MgO IBAD coatings deposited on pure Mg, and the AZ91 and AlMgSi0.5 alloys as a function of the energy and angle of incidence of the assisting Ar+-ions. The corrosion behaviour was studied by potentiodynamic
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controlled current-potential measurements under pitting corrosion conditions and the standard salt spray test. The degree of crystallinity and texture depended on the deposition parameters (i.e., I/A, ion energy, angle of incidence of the ions) and had a strong influence on the performance of the coatings. Those coatings with strong <200> texture obtained at high ion energy (10 keV) gave very poor corrosion protection due to their columnar texture and high porosity. On the contrary, highly amorphous MgO films, assisted with Ar ions at energies 3-5 keV, generally showed a good protection against corrosion. 4.2.4 Zinc and Zinc alloys Zinc coatings are well known to improve the corrosion resistance of steels by a sacrificial cathodic protection mechanism and the formation of a passive barrier that prevents further electrochemical reactions. Magnesium and zinc are commonly used as sacrificial coatings for steel in automotive, building and household applications. The common industrial processes used to deposit these coatings are galvanization and electro-deposition. These are cost-effective processes that allow high efficiencies and low cost. Alloying Zn with Al for galvanisation, and with Ni, Cr and Co for electro-deposition, have lead to significant improvements of corrosion resistance. However, the demand for increased endurance and quality coatings have driven the research to achieve better corrosion and scratch resistance. Although still at a laboratory level, PVD and IBAD techniques have been used for such purposes. IBAD Zn and Zn-Cr coatings deposited on steel have shown good adhesion and corrosion resistance. The corrosion protection was, however, observed to be thickness dependent, with thicker films presenting the best results, thus suggesting that the presence of pores is a determinant in the performance (Sansom et al., 1996; Alonso et al., 1998). Large area coil coating systems equipped with electron beam evaporation and ion beams have found little industrial acceptance, probably because the high speed at which they are commonly run. However, Wolf et al. (2000 and 2001) have recently presented a prototype at a laboratory scale for ion assisted vapour deposition in the coil coating mode. Zn alloy coatings like Zn/Ti, Zn/Cr or Zn/Al were deposited in a coil coater in vacuum by ion assisted e-beam evaporation and compared with 8 um thick electro-galvanized samples. The samples were tested according to die salt spray test. An essential improvement of the corrosion protection was obtained after pre-cleaning the substrates by ion bombardment to enhance adhesion. In addition, the results showed that the steel coated by IBAD
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185
with 4-6 |am of the different alloys provided better corrosion protection (salt spray test) than the 8 um thick electro-galvanised coatings.
4.3. Modification of biomaterials Biomaterials modification by ion-beam processing to improve the functionality and biocompatibility of some medical implants is becoming widely used. In general, it has been recognised that an adequate surface engineering of biomaterials using ion beams or plasmas leads to better and longer-lived medical implants. Recalling the characteristics of the IAD methods it is clear that ion processing can provide better bio-coatings with higher adhesive strength to the substrate. These beneficial effects are due to better control of the microstructure and chemical composition of the coatings, as compared with more conventional coating methods, e.g., plasma spraying, ion beam sputtering and non-assisted PVD methods. The area of biomaterials is probably where the expansion of the IAD methods has been more significant in the last years, probably because it is in this area where the quality is more important than the price, and where IBAD's unique characteristics are recognised and fully utilised (Cui and Luo, 1999).
4.3.1. Fretting wear and damage Fretting occurs whenever two contact surfaces suffer oscillatory movements of small amplitude for a large number of cycles. It often occurs in mechanical joints of vibrating structures in engines, trains and orthopaedic implants. Usually this phenomenon leads to both wear and fatigue damage. Since it is closely related to wear, corrosion and fatigue, the application of surface modification methods like IAD processing was seen as a way of improving the performance of biomaterials against fretting. In fact, there are some reports on the benefits of using IBAD methods to lower the fretting damage of certain coatings. The base materials most commonly used in biomedical implants are stainless steel, and different Co-. Ni- and Ti-based alloys (e.g., AISI 316L, Ti6A14V, etc.). Ti-alloys like Ti6A14V are widely used due to their high strength to weight ratio, excellent corrosion resistance and exceptional biocompatibility. However, they have a poor fretting wear and fretting fatigue resistance, two shortcomings that have to be improved for those applications (Fu et al., 1998).
186
Low ENERGY ION ASSISTED FILM GROWTH
Several coatings and surface treatments have been suggested in the literature to improve the fretting wear and fretting fatigue resistance of Ti6A14V, e.g., IBAD CrN or CuNiln coatings and surface treatments like shot penning. Fu et al. (1998) have compared these treatments. The corresponding tests showed that whereas the IBAD CrN film exhibited the best fretting fatigue performance, the duplex treatment by shot penning and IBAD CrN exhibited the highest fretting wear resistance. The authors explained the reduction of the friction coefficient in terms of a compressive residual stress induced by the ion bombardment, while the increase of hardness and surface roughness would explain the different fretting performances. By contrast, the IBAD CuNiln coatings gave the poorest fretting performance although this coating yielded the lowest friction coefficient compared with the other treatments. It appears that due to its low hardness the wear rate is the highest. High temperature (550°C) nitridation using a PHI process (Johns et al., 1996) has been shown to yield a substantial improvement in the tribological properties of the Ti6A14V. PHI increased significantly the hardness and wear resistance of the alloy, so that the wear rate was typically reduced by four orders of magnitude compared to the untreated alloy.
4.3.2. Corrosion protective coatings Different ceramic materials, such as A1203, Ti0 2 , Si0 2 and Zr0 2 are usually deposited on the base material as protective coatings against corrosion. In order to improve their adhesion, these coatings are usually deposited by IBAD. In many cases these ceramic coating are also IBAD coated with a thin (50-150 nm) silver layer to reduce the risk of bacterial infections. A comparative study of the behaviour of those ceramics showed that under simulated physiological conditions the pure and silver doped Zr0 2 coating had the best performances, i.e., highest biocompatibility and longest lasting bactericidal properties (Meinert et al., 1998).
4.3.3. Hydroxyapatite Due to its good biocompatibility and enhancement of osseous-integration, hydroxyapatite (Cai0(PO4)6(OH)2) constitutes an important coating material for
APPLICATIONS OF IAD PROCESSING
187
dental and orthopaedic implants. Although the method most frequently used to deposit hydroxyapatite (HAP) is plasma spraying, long-term clinical tests have found adhesion deficiencies, as well as high dissolution rates in aqueous solutions that limit the life of the implant. By contrast, results found in the literature (Cui et al., 1997; Ektessabi, 1997; Kim et al., 1998; Choi et al., 2000) evidence that IBAD is able to promote a higher HAP/substrate adhesive strength than plasma spraying and other conventional methods (e.g., sputtering, electron beam evaporation, laser ablation, etc.). Scratch tests performed by Ektessabi (1997) on 02+-beam assisted sputtered hydroxyapatite coatings on Ti-6A14V and steel showed an improvement of the critical load up to a factor of two with respect to the reference coatings deposited by simple ion beam sputtering (IBS). A hydroxyapatite-based coating in which calcium was partly replaced by silver has also been tested as anti-microbial. The results were similar to those reported for HAP coated with a silver film. In fact IBAD provides the most adequate silver coating on catheters and other implants with respect to the adhesion and antimicrobial behaviour.
4.3.4. Biocompatibility The relationship between surface properties and biocompatibility is a topic of interest in the research and development of biomaterials. Obviously, the surface properties of the material will determine the response of the living tissue to the implant. Corrosion and wear are two important factors to be considered in biomaterials, but tissue and blood compatibility are also required to diminish the cellular damage and blood coagulation. TU6A1-4V alloys and low temperature isotropic (LTI) pyrolitic carbon are widely used in biomedical applications because of their high corrosion resistance and biocompatibility. Owing to these properties, these materials are used in the fabrication of heart valves, hip joints, etc. However, several coatings and surface modifications are also employed to improve their wear and corrosion resistance as well as their biocompatibility. In fact thrombogenic problems, which have to be treated with anti-coagulants, are rather common in patients with implanted heart valves. Wang et al. (2000) have studied the biocompatibility of TiOx films deposited on LTI by IBAD. These authors claim that these coatings improve the blood compatibility of the LTI, both in-vivo and in-vitro tests. The behaviour is
188
Low ENERGY ION ASSISTED FILM GROWTH
explained as due to the presence of Ti and Ti oxidation states at the surface of the coating, which make the surface more polar and reduce the work function (Wang et al., 2000). DLC and CNX films have also been proposed as candidates for this application because their chemical inertness and composition make them biologically compatible. Cui and Li (2000) have reviewed the biocompatibility of DLC and CNX films. In general, the results are rather promising. Both materials, as deposited by IBAD methods, show good adherence, chemical inertness and good biocompatibility towards various cell types.
4.4. Metallisation of polymers IAD processes are highly suitable for coating polymeric substrates where deposition must be done at low temperatures. Many industrial applications of polymers require the deposition of thin metallic films to modify their surface functional properties like electromagnetic shielding, wear protection, gas diffusion barrier, etc. However, the adhesion of inorganic films on the smooth and chemically inert surfaces of polymers is usually poor. In general, improvements of adhesion to levels of peel strength above lN/mm are necessary for many applications and therefore the metal-polymer interface becomes a crucial parameter. The use of IAD methods to improve the adhesion between dissimilar materials was discussed in section 3.11. Specific preconditioning of polymers to make surface coatings ready for industrial applications is usually necessary. The most commonly used preconditioning process is the chemical attack with chromic acid. However, the social demand for clean technologies has stimulated the development of ion beam and plasma assisted processes for the conditioning and metallisation of polymers. Kupfer and Wolf (2000) have shown that ion beam preconditioning with Ar-ions of different energies and fluences leads to the formation of graphitic carbon on the surface of the untreated Poly(phenylene sulfide) (PPS) and hinders the adhesion of the deposited metal. On the contrary, ion beam assisted evaporation of Cu leads to a significant improvement (up to 2.5 times) of adhesion on PPS. Figure 4.3 shows the pull-off strength of IBAD Cu coatings on PPS as a function of the I/A ratio and energy of assisting ions. Interestingly, the improvement of adhesion is only obtained at low energies and low I/A ratios. If the input energy of the assisting
189
APPLICATIONS OF IAD PROCESSING
ions is too high the bonding at the metal/polymer interface becomes weak and the adhesion poor. The main effect of the IBAD process seems to be the formation of chemical bonds at the interface (cf., section 3.11), without affecting the roughness of the substrate. However, an excess of ion bombardment can induce a graphitisation of the interface and therefore a decrease in adhesion.
£
0.02
0.04
0.06
0.08
0.1
0.12
Ion/atom ratio Figure 4.3. Pull off strength of Cu deposited by IBAD on PPS as a function of the I/A ratio for different energies of the Ar-assisting ions. Reproduce from Kupfer and Wolf (2000).
Cu/Teflon and Fe/Teflon are two typical examples which show a negligible adhesion when deposited without the assistance of ion bombardment, either as preconditioning or during the deposition. Low energy pre-sputtering of Teflon is very effective in producing a high strength bond. Only a few seconds of ion bombardment produces a maximum in the adhesion that has been associated with the formation of C-Cu and Fe-C bonds and a graded layer at the interface (Chang et al., 1987). Nevertheless, these results are strongly dependent on the system under study and cannot be generalised. In fact, there are also examples where IBAD treatments have failed to improve the adhesion. As an example of different responses of systems prepared under the same IBAD conditions, we can use the results reported by Loh et al. (1988). These authors have studied the IBAD deposition (assistance with 400 eV Ar-ions) of Cu, Ag and Au films on Plexiglas, Teflon, Kapton and Lexan. As compared with the evaporated films, which do not pass the common scotch tape test on Teflon, but adhere well on Lexan and Kapton and fail on Plexiglas, the coatings deposited by IBAD improved significantly their
190
Low ENERGY ION ASSISTED FILM GROWTH
adhesion on Teflon, did not show any significant improvement on Lexan and Kapton and failed to achieve any adhesion on Plexiglas. Therefore, it can be concluded that each system requires an appropriate and specific treatment to enhance specific mechanisms for adhesion improvement (cf., section 3.11).
4.5. Optical coatings The fabrication of optical coatings remains the area where the use of IAD techniques, e.g., ion plating, DIBS, IBAD, etc., is more extended. Many optical coating companies have incorporated ion beam assisted processes since they experienced that ion bombardment was essential for obtaining properties that are required for several applications of the evaporated films. Some companies are nowadays using these processes for production of optical coatings which are sensitive to the atmospheric moisture or which require low deposition temperatures like heavy metal fluorides. Ion bombardment is being used for pre-deposition conditioning of surfaces of germanium, silicon and other materials which are somewhat resistant to thin film adherence (cf., section 3.11). In addition, IAD has also demonstrated its ability to increase the packing density (defined as the ratio between the volume of the solid part of the film and the total volume of the film), improve the stability and durability, modify the stress and adjust the stoichiometry using reactive ions. Ion assisted coatings are commonly associated with coatings of higher refractive index, free of spectral shifts upon exposure to atmospheric moisture, free of adsorbed water and the corresponding infrared absorption and good adhesion and mechanical properties. Nevertheless, whereas IAD methods have proven highly advantageous for the deposition of high quality dielectric films, in the case of metals and semiconductors, other techniques have resulted in films with similar or better properties. For many semiconductors, the bombardment with energetic particles usually leads to damaged lattices and poorer properties. It is interesting to observe the lack of information regarding the optical behaviour of ZnS obtained by IAD, even though it is the most widely used coating for IR applications. In contrast, the number of reports in the open literature dedicated to the optical behaviour of silica, titania, tantalum pent-oxide and many other oxides deposited by IAD is huge. Another interesting problem in the field of optical coatings is that of the laser damage. Even though it is a very important topic and some progress has been
APPLICATIONS OF IAD PROCESSING
191
made in recent years, the improvement of laser damage thresholds in thin films has not yet been solved. Furthermore, although many attempts have been made using the whole variety of IAD methods, it has not yet been concluded whether the IAD coatings give higher threshold values than the more conventional methods. In fact, it has been found sometimes (Alvisi et al., 1999) that films deposited by non-assisted methods, with lower packing density may have a higher laser damage threshold and better heat dissipation. For details on the influence of the structure of oxide coatings on its laser damage threshold the reader is referred to the review published by Hacker et al. (1996). As a rule, the laser damage threshold always remains well below that of the respective bulk materials. If the well-established benefits of IAD could somehow be associated with a high laser damage threshold, then the process would probably become universally accepted. Excellent reviews of the status of ion assisted techniques (including plasma processes) for optical thin films have been written by Martin (1986), Martin and Netterfield (1986), Gibson (1987), Mohan et al. (1995), Bovard, (1996) and many others. They should be consulted for a more detailed study of different optical materials and effects.
4.5.1. Dielectric oxide films Dielectric oxide films for optical applications are usually deposited by thermal and e-beam evaporation of oxide materials. However, as for many other evaporated films, optical thin films present a series of problems associated with the characteristics of that deposition method, i.e., loosely packed columnar structure, absorption of water, variable optical parameters, etc.. Heating the substrates up to several hundreds °C improves the density and the optical properties, but many times heating is not possible or causes other undesirable effects in the microstructure and morphology of the substrate and film. Therefore, ion beam and plasma assisted processes have found a rapid acceptance for the production of environmental stable dielectric layers and optical filters. The number of reports on the dependence of the optical properties of different dielectric films deposited by IAD methods is very large. Many of the results obtained up to the mid and late 80s have been revised in detail in the reviews by Martin (1986) and Gibson (1987). We will focus here on some general results which demonstrate some of the advantages of using IAD.
192
Low ENERGY ION ASSISTED FILM GROWTH
Figure 4.4 demonstrates the improved environmental stability achieved by using ion beams during deposition of Si0 2 films. The figure shows the evolution of the refractive index of Si0 2 coatings deposited by ion assisted deposition (Ar+ + 0 2 + ) as a function of ion energy per deposited molecule (SiC>2) (Souche et al., 1998). For each coating the refractive index was measured in vacuum and after aging in air. Figure 4.4 shows that ion bombardment induces a densification of the film and the suppression of the porous columnar microstructure shown by those films whose deposition has not been assisted or assisted with very low energy ions (i.e., Ed < 40 eV). The effect is evidenced by the shift of the refractive index shown by the poorly assisted films after aging, which is associated with the uptake of water that normally occurs in porous films upon exposure to the atmosphere. Whereas the films deposited without assistance (Ed=0) or assisted at low energies show a significant increase of the refractive index when exposed to air, the films deposited at energies above 40 eV are dense coatings which do not take in water and have stable refractive index. Similar shifts in the refractive index when the sample is transferred from the vacuum to air have been also reported for other evaporated oxide films (Martin et al., 1986).
. 1.50-
•
Ec=150eV
•
Ec=300eV
1.44-
1.41-
•
*
S4
C § d c
•
1
1.47-
S2 S1
6 1.38100 200 Average deposited energy Ed (eV)
300
Figure 4.4. Dependence of the refractive index of ion assisted Si02 films as a function of the normalised energy per molecule Ed. The films were assisted with ions with two kinetic energies Ec. The open points represent the refractive index after aging in air. Reproduced from (Souche et al., 1998) with permission.
193
APPLICATIONS OF IAD PROCESSING
Figure 4.4 also shows that the refractive index increases rather steeply as the normalised ion energy (Ec x I/M, where M is the net flux of deposited Si0 2 molecules) is increased up to a critical normalised energy of around 40-50 eV. This increase in the refractive index is observed as an increase of density and elimination of the columnar porous structure (cf., section 3.5.2). Beyond that critical value, the refractive index saturates up to 300 eV, which is the maximum ion energy reached in these experiments (Souche et al., 1998). A direct comparison between the dispersion characteristics of S1O2 and Ti0 2 films deposited by IAD and electron-beam evaporation (EBD) as published by Tsai et al. (1997) is shown in Fig. 4.5. This figure clearly shows that the films prepared by reactive IAD have higher refractive indices than those deposited by reactive EBD under similar deposition conditions.
1.49
— 1.47 '.oo. 0
400
Si0 2 ,EBD
••<>••• Si0 2 ,IAI>
500
600
700
800
>
— Ti0 2 ,EBD TiQ2,IAD O-OOOOOOO-o-OOO-Oo I
400
500
600
700
800
Wavelength (nm) Figure 4.5. Dispersion curves of Si02 and TIO2 films deposited by reactive IAD and electron beam evaporation (EBD). Reproduced from Tsai et al. (1997) with permission.
194
Low ENERGY ION ASSISTED FILM GROWTH
The above described dependence of the refractive index on the assisting ion energy and current density has been observed for many other oxide films (Mohan and Krishna, 1995) and interpreted as evidence that the film density and refractive index could be increased up to values close to the respective bulk values. In many cases, at high energies of the assisting ions, the refractive index is observed to fall due to an excessive damage, preferential sputtering of oxygen or significant incorporation of the assisting ions into the growing film. The existence of a critical energy and I/A ratio, beyond which the refractive index saturates or falls down was clearly stated by Martin et al. (1986) for oxide films deposited by IBAD methods. Although the critical value of the energy for a given I/A ratio and vice versa are obviously dependent on the material, Martin et al. (1986) suggested that assisting ions with energies 300-600 eV and ion/atom arrival rate ratio (I/A) between 0.1 and 0.3 (for ion current densities in the range 200-250 uA cm"2) were ideal for producing oxide films with maximum refractive indices. In some cases, (e.g., Ce0 2 ) the use of energetic oxygen ions allows stoichiometric oxide coatings even at low deposition temperatures (Netterfield et al., 1985). The increase of the refraction index of oxide films deposited by IAD methods has been confirmed in all the oxides investigated using different deposition techniques under adequate deposition conditions. A compilation of values of the refractive indices at 550nm found in the literature for both IAD and EBD deposited dielectric materials is presented in Table 4.1, showing rather clearly that the refractive indices of the IAD films are higher than those for evaporated films. The increase in the refractive index is attributed to an increase in packing density, which is a result of the ion assistance. In general, while evaporated films show a packing density of 0.7-0.8 at room temperature up to 300°C, IAD films can reach packing densities close to 1 (i.e., bulk) depending on the assisting conditions. Using the well known Maxwell-Garnett or Bruggeman relationships within the effective medium approximation (Tompkins, 1993), the pore volume fraction and water filling ratio can be estimated from ellipsometric measurements. In fact, it is rather common to estimate the packing density or the void volume fraction in terms of the refractive index of the film. Figure 4.6 shows the variation of the pore volume fraction of SiC>2 films deposited by ion assisted e-beam evaporation as a function of the energy deposited per Si0 2 molecule Ej as described above. The data have been published by Brunet-Bruneau et al. (1998) and indicate that upon increasing the energy of assistance Ed the volume fraction of voids is significantly reduced until a constant residual volume of 0.05 is reached for Ed >100 eV. A
195
APPLICATIONS OF IAD PROCESSING
comparison of this behaviour with that shown in Figure 4.4 for the refraction index clearly confirms that for Ed <100 eV the main effect of the assisting ions is to increase the density of the silica films. Table 4.1. Range of reported refractive indices at 550 nm (after air exposure) of oxide films obtained by IAD methods as compared with evaporated films under different deposition conditions.
refractive indices of EBD films
refractive indices of IAD films
Si0 2
1.46
1.46-1.49
McNeil etal. (1984), Pawlewicz et al. (1994), Tsaietal. (1997), Souche et al. (1998), Niederwald et al. (1999)
A1203
1.54
1.67-1.7
McNeil etal. (1985), Pawlewicz et al. (1994)
Ti0 2
1.90
2.39 - 2.52
Martin etal. (1983), McNeil et al. (1985), Pawlewicz et al. (1994), Tsai et al. (1997)
Ta2Os
1.99
2.15-2.18
Martin et al. (1983), Pawlewicz et al. (1994), Cevro et al. (1995)
Zr0 2
1.93
2.05 - 2.21
Martin etal. (1983), Krishna etal. (1992), Pawlewicz et al. (1994)
Ce0 2
1.95
2.3 - 2.49
Netterfield et al. (1985), Al-Robaeeetal. (1991)
material
references
The optical losses of the oxide coatings play an important role in their performance also and therefore they must be considered. In general, the extinction coefficient of the oxide films deposited by IAD methods has been observed to be below the tolerable limits (i.e., <10"3) and in any case lower than the values of films deposited by evaporation. The absorption of high refraction index oxides, e.g., Ti0 2 Ta 2 0 5 , etc. has been found to be significantly reduced by IAD processes. However, if the assisting parameters (e.g., energy and current density or I/A ratio) go beyond certain values which depend on the material, the absorption also increases. The
196
Low ENERGY ION ASSISTED FILM GROWTH
d volume fraction
effect has been reported for all the dielectric oxides Ta 2 0 5 , A1203, Ce0 2 , Ti0 2) Zr0 2 and has usually been attributed to a significant damage of the film, a loss of stoichiometry due to preferential sputtering of oxygen, or to the incorporation of the bombarding species (e.g., 0 2 , Ar) in the growing film.
( > Sample A 0.200.150.10-
•1 •2
••
9
•
0.05
•mm
mu'
•
A Ail
o.uo
()
100
200
300
Ed (eV) Figure 4.6. Variation of the void volume fraction of Si02 films deposited by IB AD as a function of the normalised assisting energy per deposited molecule Ej. Reproduced from Brunet-Bruneau et al. (1998) with permission.
Surface roughness as well as film microstructure are also important factors directly related to the optical scatter and the coating performance for most applications. In general, ion bombardment during film growth has shown to reduce surface roughness (cf., section 3.3). Al-Jumaily et al. (1985) have compared the scatter spectrum of Si0 2 and Ti0 2 films prepared by evaporation and IBAD. In both oxides the IBAD films resulted in a smoother surface and in a much less optical scatter characteristics.
4.5.2. Fluoride thin films Thin films of fluoride materials (e.g., A1F3> MgF2, ThF4, BaF2, CaF2 etc.) are widely used for optical interference coatings applications. The attractive features include wide optical transparency range, low index of refraction, high laser damage threshold and low dispersion. Most fluorides transmit from the UV to the mid-IR. All those properties make metal fluorides very useful as low index film in bandpass interference filters from the UV to the mid-IR. Unfortunately, evaporated films are
APPLICATIONS OF IAD PROCESSING
197
characterised by poor optical, chemical and physical properties. The usual columnar structure with packing densities of about 60-80% yields highly stressed films with poor and unstable optical properties. All that makes IAD methods very appropriate for the deposition of fluoride coatings with a good optical behaviour. IAD improves the compactness and reduces the tensile stress of the films. Although Ar+ is commonly used as assisting ion, mixtures of CF4 and 0 2 are also incorporated to keep the stoichiometry and to reduce the absorption of the film in the UV range . Illustrations of the reduction of water absorption in fluoride thin films deposited by IAD methods have been reported very often. Typical results indicate that when the films are exposed to the atmosphere after deposition, water vapour is incorporated in the films in decreasing amounts as the ion assistance of the deposition increases. The inclusion of water vapour in the film is shown by the presence of a strong IR absorption band at 3 urn (Al-Jumaily et al., 1987; Scaglione et al., 1992). In summary, the main advantages of using IAD for optical coatings are, improved optical density leading to higher index of refraction, lower absorption and scatter and improved environmental stability. Moderate levels of ion bombardment lead to stress free coatings or moderate levels of compressive stress which enhances adhesion. In fact, stress is a major problem in filters with a large number of layers and applications where the required film thickness is high (e.g., IR applications) (cf., section 4.5.3). It has been shown that one can remedy many of the problems of the evaporated films by raising the substrate temperature. However, in some applications where the substrate is highly sensitive to the temperature, e.g., III-V or II-VI semiconductors, heavy metal fluorides, plastics, etc., IAD at low temperatures are the most appropriate methods. Usual problems associated with the IAD techniques are, preferential sputtering of any of the components leading to problems of stoichiometry and/or significant damage of the film when the ion bombardment is too intense, an effect that increases the absorption of light and decreases the transparency of the films.
4.5.3. Narrow band filters Optical interference filters are made by depositing a sequence of discrete layers of transparent materials that have different refractive indices. Dielectric films having a band-gap beyond the visible are used extensively for multilayer wavelength-
198
Low ENERGY ION ASSISTED FILM GROWTH
selective coatings. Of these materials, those combinations of especially high and low refractive index (e.g., Ti0 2 /Si0 2 , Ta2CVSi02, MgF2/Zr02 etc.) are particularly appropriate, since the index difference between adjacent layers determines the number of quarter waves required to achieve effective antireflection or reflection enhancement. Obviously, the stability of the centre wavelength is particularly important for these filters with a narrow band-pass. In fact, it is a well known problem that once a filter is produced by conventional methods (e.g., e-beam evaporation), removed from the vacuum chamber and exposed to the atmospheric water vapour causes the centre of the filter to displace towards higher wavelengths (Takashashi, 1995). This effect is explained by the fact that the effective refraction index of the film varies according to the level of absorbed water. As the amount of absorbed water is dependent on the atmosphere humidity and temperature, the filter characteristics become dependent on these two factors. However, as demonstrated above (cf., section 4.4.1), the use of IAD techniques makes it possible to grow dense films of near unity packing densities and therefore less affected by the ambient moisture and temperature. The improvement which supposes using IAD methods to deposit UV-IR cutoff filters has been nicely demonstrated by Tsai et al. (1998). These authors have compared Ti0 2 /Si0 2 filters produced by IAD (at room temperature) and electron beam deposition (EBD) (at 300°C) with respect to the band width and thermal stability. The results presented in Figure 4.7 correspond to UV-IR cutoff filters consisting of alternating Ti0 2 and Si0 2 layers to form a superposition of a 24-layer near-IR cutoff filter and a 18-layer UV cutoff filter on the first and second surfaces of the glass substrate. Figure 4.7 a) shows the spectral transmittance of the filters prepared by IAD and EBD (Tsai et al., 1998). The figure shows that the bandwidths (UV and near-IR blocking edge) for the IAD filter are wider than for the EBD filter. Since both filters have the same number of layers and optical thickness, this difference is clearly a consequence of a higher ratio of refraction indices, n(Ti02)/n(Si02) for the films deposited by IAD (n(Ti02) = 2.386, n(Si0 2 ) = 1.461) as compared with those produced by EBD (n(Ti02) = 2.167, n(Si02) = 1.446) (cf., Figure 4.5). Figure 4.7 b) shows the near-IR cut-off edge position of the filters deposited by IAD and EBD as a function of the temperature. In this case, we observe that the temperature stability of the IAD filters is superior to that of the EBD filters. This all confirms again that the IAD coatings are denser and more stable than those deposited by evaporation.
199
APPLICATIONS OF IAD PROCESSING
S e o o o. 0)
IAD
M
EBD
•o SB e •
•4->
a u
300
500
700
Wavelength (nm)
900
I
" 0
1 50
'
1 ' 1 ' 1 ' 1 100 150 200 250
r
300
Temperature (°C)
Figure 4.7. a) Spectral transmittance of two near-IR cutoff filters prepared by IAD and EBD methods and b) behaviour of the near-IR cut-off edge as a function of the temperature. Reproduced from Tsai et al. (1998) with permission.
Another advantage of using IAD methods for multilayer filters and thick films, such as those required for infrared applications, is its ability to control stress levels (cf., section 3.10.5). In many cases, the stress can be the limiting factor in determining the maximum usable thickness without leading to delaminating.
4.5.4. Rugate filters One of the most significant advances in optical filters has been the realisation of rugate type filters. A rugate filter is an interference filter in which the refractive index varies periodically in a smooth and continuous manner. The simplest example of that is when the variation is sinusoidal. The main advantage of the rugate with respect to the multilayer filters is the lack of interfaces that cause stress and delaminating as well as light scattering. Its realisation requires a material whose refractive index can be varied significantly as a function of the composition and a deposition technique which enables a control of the composition (refractive index) in terms of the deposition parameters.
200
Low ENERGY ION ASSISTED FILM GROWTH
The first report of a continuous modulation of composition by IBAD was done by Donovan et al. (1989), who investigated SiNx films as a function of the nitrogen ion current density. The refraction index was observed to vary continuously from n=3.9 for Si up to n=2.0 for Si3N4 by suitable adjustment of the Si to N flux ratios. These authors also realised a rugate filter formed by 23 cycles of a sinusoidally variable refraction index (i.e., composition), whose experimental transmission was in good agreement with the simulated one. Ion assisted deposition of SiNx films performed by e-beam evaporation of silicon and bombardment with only nitrogen ions has also been reported by Lee et al. (1999). The SiNx films were all amorphous and the refractive index ranged from 1.72 to 3.43 depending on the current density for an optimum energy of the nitrogen ions of 550 eV. Above this energy, the extinction coefficient increases significantly. Their application as near IR antireflection coating and as a band pass filter based on multilayers of Si and SiNx was demonstrated. However, the k values were relatively high (~102) for x<1.33, which constitutes a limitation if a low loss coating is required. Silicon oxynitride (SiOxNy) deposited by IAD is also a potential candidate in rugate filters because the refractive index can be varied gradually from that of Si0 2 to that of Si3N4. This possibility has been investigated by Cho et al. (1997). These authors have investigated the optical properties of inhomogeneous silicon oxynitride films prepared by nitrogen assisted e-beam evaporation of silicon and demonstrated the possibility of obtaining variable refractive indices as a function of the composition. To control the composition of SiOxNy either the oxygen pressure was varied at a fixed ion energy (500 eV) or the energy of the nitrogen ions was varied without any backfilling of oxygen. Figure 4.8 a) shows data of Cho et al. (1997), where the refractive index and oxygen content of SiOxNy films have been depicted as a function of the oxygen backfill pressure under constant nitrogen bombardment (500 eV and 28 uA/cm2). The refractive index varies continuously between 2.02 for Si3N4 and 1.49 for Si0 2 , while the extinction coefficient was always ~10"3. Figure 4.8 b) shows the variation of the refractive index and oxygen composition as a function of the nitrogen ion energy with no oxygen backfilling. In this case, the refractive index decreases from 2.02 at 500 eV up to 1.82 at 700 eV and then increases up to 2.27 for 1000 eV. However, the range of low absorbing films is only between 1.82 and 2.02 at energies between 500 and 900 eV. As the energy increases above 900 eV the
201
APPLICATIONS OF IAD PROCESSING
absorption increases significantly as in the case of dielectric films (cf., section 4.5.1). Using the first method, the authors realized a simple rugate filter whose transmittance agreed rather well with that simulated for a continuously varied sinusoidal index. The variations in the refractive index are smaller than for SiNx but the optical losses are also lower.
Oxygen pressure (Xlff Torr) 2.3 • •
•
2.2'
index Qz comp.
•
SH>2
•
| , 1 |
2.,
I
•
I*S 1.9 1* 1
500
600
1 —r — i •
800
900
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1000
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Ion Energy (eV) Figure 4.8. Variation of the refractive index and oxygen composition of SiOxNy films as a function a) the oxygen backfill pressure for assisting nitrogen ions of 500 eV (28 uA/cm2 ), b) as a function of the nitrogen ion energy without oxygen backfill. Reproduced from Cho et al. (1997), with permission.
4.5.5. Transparent conducting films Optoelectronic devices require transparent electrical contacts to allow charge carriers to be injected into or extracted from active semiconductor material while allowing light to enter or exit the device (LED, photovoltaic cells, photodetectors,
202
Low ENERGY ION ASSISTED FILM GROWTH
etc.). Although a large variety of transparent and electrically conductive oxides (TCO) are known, the two most widely used TCO's are indium tin oxide (ITO) and aluminium zinc oxide (ZnO:Al, AZO). In fact, ITO is the most extensively used because of its relatively low resistivity (< 10"5 Q m) and high transparency in the visible, above 80-85% at 550 nm. ITO has found numerous applications as transparent electrodes of flat panel displays and in other optoelectronic and photovoltaic devices. However, these technologies very often demand the use of temperature sensitive substrates (e.g., organic materials), which also require special deposition techniques. In particular the IAD methods seem very suitable for a successful room temperature deposition of ITO films, so that a large variety of IAD techniques and deposition conditions have been investigated to further improve the conductivity and other properties of ITO films deposited at low temperature. In general, deposition at room temperature leads to transparent films with a high refractive index. However, in order to obtain low resistivity values (e.g., 10"6 Q. m) higher deposition temperatures or treatments after deposition are required. Transparent ITO coatings prepared by biased magnetron sputtering have been reported in the literature (Yang et al., 2000) and are also available commercially with resistivities as low as 2xl0"6 £2 m. The use of unbalanced planar magnetron sputtering and ion-assisted plasma anodisation to provide Ar+-bombardment has been reported by Danson et al. (1998). When the deposition is performed at room temperature these authors report ITO films with a transmittance of over 80% and a resistivity around 10"5 Qm. Amorphous ITO films deposited by IBAD with a mixture of 0 2 and Kr ions at room temperature with good resistivity (4xl0 6 Qm) and transmittance (90% at 550 nm) have been reported recently by Kim et al. (2000). Kim et al. (2000) have reported on the influence of Kr+ bombardment during growth of ITO thin films deposited on polycarbonate and glass at room temperature by 0 2 + assisted evaporation. Whereas the evaporation rate of bulk ITO and the flux and energy of the oxygen ions was kept constant (the flux and energy of the assisting Kr+ ions was varied to study their effect. The results indicate that both the resistivity and the optical transmittance are improved by increasing Kr+ bombardment if the energy is maintained below 400 eV. Figure 4.9 shows the resistivity and the transmittance at 550 nm of the ITO films deposited on both substrates (i.e., glass and polycarbonate) at room temperature as a function of the flux of 200 eV assisting Kr+ ions.
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APPLICATIONS OF IAD PROCESSING
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The results indicate an improvement of the packing density and film stoichiometry by extra Kr+ bombardment. However, increasing the energy above 400 eV leads to a significant decrease of the transmittance, probably due to damage and increased surface roughness. In general, it is found that the energy of the incident ions or species is a key deposition parameter to reduce the resistivity of ITO films. It is preferable to use higher ion flux at lower energy than the contrary, but if the energy is too low the oxidation process and, therefore, the stoichiometry of the film could be affected. A process based on the use of oxygen cluster ion beams (i.e., GCIB) to assist the simultaneous evaporation of In and Sn seems to meet the appropriate characteristics to obtain high quality stoichiometric ITO films at room temperature. The method has been applied by Qin et al. (1998) and Matsuo et al. (2000) to obtain
204
Low ENERGY ION ASSISTED FILM GROWTH
very smooth, highly transparent (>80%) and low resistivity (~ 5xl0"6 D. m) ITO films using oxygen cluster ion beams of 5-7 keV. The formation of the stoichiometric oxide was observed to depend on the acceleration voltage, i.e., the kinetic energy of the clusters. For an oxygen cluster mean size of 2000 molecules the formation of In203 requires energies above 5 keV (ion current density of 100 nA cm'2), as it is shown in Fig. 4.10. It is interesting to observe that a cluster of 2000 molecules with an energy of 5 keV corresponds to an average kinetic energy of 2.5 eV per molecule, well below the threshold to damage the film but in any case necessary to form stoichiometric films.
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Acceleration voltage Va (kV) Figure 4.10. WO ratio as a function of the acceleration voltage of oxygen ion clusters (average size 2000 molecules) assisting the deposition of ITO films. Reproduced from Matsuo et al. (2000) with permission.
4.6. Magnetic thin films Tremendous progress has been made in the field of magnetic materials research and technology over the last few years. Better properties and novel structures have arisen due to the ability to synthesise new structures or to tailor microstructures. Processing by IAD methods modifies the microstructure, e.g., surface and interface roughness, interface mixing, texture, stress, etc. and therefore can influence magnetic properties like magnetic anisotropy, coercivity, magnetoresistance, etc. Ion beam deposition (IBD) has found a wide acceptance in the field of
APPLICATIONS OF IAD PROCESSING
205
writing/reading magnetic heads. IBD allows the control of both the energy and angular distribution of the deposited species upon arrival at the surface and therefore enables a better control of the smoothness and mixing at the interfaces, which are very important characteristics by the deposition of GMR multilayers.
4.6.1. Thin metallic films Many different deposition techniques are used to apply magnetic thin films for recording, including PVD, magnetron sputtering and, obviously, IAD methods. Although the number of works reporting the influence of the deposition parameters on the magnetic properties is rather scarce, we present in the following some examples where the deposition method and parameters were correlated with the magnetic properties of the deposited metal. Ultrahigh purity Fe and Ni are of particular interest for thin film (soft) magnetic heads. The main problem of these films, as well as that of many other metallic films, is their growth with a spongy columnar structure which favours their oxidation and the instability of their properties. The beneficial influence of the ion assistance on the magnetic properties of Fe thin films deposited by DIBS have been reported by Nagakubo et al. (1989). These authors showed that in comparison with the non-assisted films, whose magnetic properties were highly dependent on the residual pressure in the deposition chamber, the Ar+-assisted films showed stable and reliable saturation magnetisation Ms and coercive field Hc. The better and more stable properties of the films deposited by DIBS were associated with the loss of the columnar structure and the densification of the films. The structure and the electrical and magnetic properties of nickel films obtained by IBAD has been reviewed by Ensinger (1998). According to this review, the magnetic properties of Ni films appear to be strongly dependent on the I/A ratio during deposition. The coercive magnetic field Hc has been observed to decrease with the I/A ratio for low ion energies due to an enhancement of the micro-structural order and purity (less oxidised) of the films. However, if the beam energy is increased, so that ions can induce enough structural disorder and damage to overbalance the reduction of impurity content, an increase of Hc. is observed Another interesting effect that can be induced by ion bombardment is the anisotropy of the magnetic properties. This effect has been demonstrated by Lewis
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Low ENERGY ION ASSISTED FILM GROWTH
et al. (1992), who deposited Ni films on quartz by DIBS using Xe ions for both sputtering and assistance. When the growing films were bombarded with 0.1 keV Xe+ at 45°, the films developed an in-plane uniaxial magnetic anisotropy with the easy axis laying perpendicular to the plane of ion incidence. However, those Ni films deposited without ion irradiation were isotropic. The degree of anisotropy depended on the assisting conditions. Lewis et al. (1992) conclude that the cause of this magnetic anisotropy is the development of strain perpendicular to the film plane and the corresponding inverse magnetorestrictive effect. The magnetic properties of IAD films can be influenced by the ion bombardment due to both the modification of the microstructure and the stress state of the films. For metals in general, a moderate ion bombardment with low ion energies usually enhances the crystallinity of the metal and induces a (111) texture (cf., section 3.9.3). Off-normal incident ions induce in-plane grain orientation and therefore magnetic anisotropy.
4.6.2. Magnetoresistive materials Magnetoresistive (MR) materials are characterised by their property to display a change in their resistance when exposed to a magnetic field. Giant magnetoresistance refers to large changes in resistivity obtained with modest fields (< 10 3 Oe) in carefully designed metallic films consisting of ferromagnetic and nonmagnetic components. Since the discovery of the Giant Magneto-Resistive (GMR) effect in multilayers consisting of the stack of successive antiferromagnegtic (AF) and ferromagnetic (F) layers, investigations of exchange coupling between AF and F bilayers and their dependence on the deposition methods and parameters have been carried out by many researchers because of their applications in MR spin valve sensors (Kools, J. C. S., 1995). However, the number of examples where, the magnetic properties of MR materials and GMR systems, deposited by IAD methods, have been addressed is rather scarce. Actually, although it is generally accepted that a sharp interface contributes, to a large extent, to a better coupling between the F and AF films, other factors like texture, impurities, etc., have to also be considered. Since the discovery of the GMR properties, different multilayer systems (e.g., NigiFeig/Cu, Co/Cu, etc.) have attracted considerable attention.
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207
The usual deposition technology of these multilayers is magnetron sputtering and the use of IAD methods was in question due to the risk of inducing ion beam mixing at the interfaces. Nevertheless, there are some reports on the possibilities of IB AD to engineer the coercivities of the ferromagnetic layer. Schmeusser et al. (1997) have reported on the deposition of Co/Cu multilayer films by ion beam sputtering (IBS) and magnetron sputtering (MS), varying the deposition conditions to optimise the GMR effect. The goal was to find out whether the films deposited by IBS could compete with those obtained by plasma sputtering, which usually present larger magnetoresistance. In general, two main deficiencies of the IBS method relative to the MS were identified. One is the much lower deposition rate which leads to a higher incorporation of impurities as demonstrated by a larger saturation resistivity. The other deficiency is the interface mixing induced by the high energy particles. Both effects could be reduced by replacing Xe for Ar (i.e., increasing the ion mass) as sputtering gas, using the lowest ion energy and positioning the substrate far away from the forward sputtering direction. The use of DIBS for controlling coercivities in NigiFeig/Cu multilayers has been demonstrated by Gutierrez et al. (1997). These authors were able to obtain multilayers with low resistivities (2x10"7 Q m"1) and nearly bulk like saturation magnetization. The coercivity of the moderately ion-assisted films was observed to increase with respect to the non-assisted. However, a more severe ion bombardment induces ion beam mixing at the interface and a drastic decrease of the magnetization of the Ni81Fe!9 layer. Therefore, to optimise the system, the ion assistance was suppressed during the initial deposition of Ni8iFei9 to avoid the mixing effect at the interface, while it was then applied during the growth of the film to obtain the maximum coercivity, all of which demonstrates the feasibility of using low energy ion beams to selectively manipulate the coercivity of the components of a multilayer system. Lee et al. (1998) investigated the exchange field of NiO/NiFe bilayers deposited by DIBS as a function of the deposition parameters. Both the effect of the current density and energy of the sputtering beam and ion assistance of the growing NiO film on the exchange bias and coercivity of NiO/NiFe bilayers was investigated. The results indicate that ion beam sputtering (without assistance) is able to produce a smooth fine grained NiO surface which leads to a large exchange field of about 100 Oe. These authors do not observe any enhancement of the
208
Low ENERGY ION ASSISTED FILM GROWTH
exchange coupling by the use of a secondary ion beam to bombard the growing NiO film, but rather an increase of the interface roughness and a change in the texture of the NiO from (220) to (200) as the energy of the Ar+ was varied between 500 and 1500 eV. An interesting application of IBAD to obtain an enhancement of the exchange field in sputtered NiFeO/NiFe bilayers due to a better in-plane orientation of the NiFeO layer has been demonstrated by Lai et al. (1999). These authors use IBAD to produce in-plane textured MgO films (cf., Section 3.9) which are then used as structural templates to modify the crystalline orientation of NiFeO. The exchange field was maximum when the in-plane orientation of NiFeO was aligned more perfectly.
4.6.3. Reading/writing magnetic heads In the last decade there has been a large interest in the development of multilayerstructures consisting of alternating F and AF layers which show GMR effect and the corresponding deposition technologies. Kools et al. (2000) have reviewed both the materials and vacuum deposition technologies utilised for the fabrication of reading magnetic heads based on GMR multilayers. According to this review, the most common design of a read head consists of a stripe of magnetoresistive material placed in between two soft magnetic fields. The spin-valve, or more generally GMR is a series of metallic multilayers consisting of a large number of individual layers (up to 20) of ferromagnetic (Ni81Fel9, Ni78Fel9Cr3, Co90Fel0), antiferromagnetic (e.g., NiO, Ir20Mn80, Pt48Mn52) and non-magnetic materials (e.g., Cu, Ru, Ta). In general, the films are typically 2-3 nm thick, polycrystalline, with commonly (111) oriented grains ( Kools et al., 2000). Obviously, the microstructure will depend on the deposition method and parameters (cf., Chapter 3). In general, it is desirable to have large and oriented grains, as well as abrupt and smooth interfaces, because both features result in better magnetic, electric and magnetoresistive properties. In addition, contamination and collisional mixing due to energetic particle bombardment have to be avoided. These requirements make the choice of a deposition method very critical, as it has to offer complete control of the microstructure and morphology through the correct choice of the deposition parameters. In fact, different deposition methods (e.g., IBD, MS, etc.) have been applied and optimised, but apparently IBD seems to be the appropriate choice (Kools et al., 2000). IBD allows the control of the energy and angular distribution
APPLICATIONS OF IAD PROCESSING
209
of the deposited metal atoms upon arrival at the surface and therefore enables to control the smoothness and mixing extent at the interfaces. To avoid mixing at the interfaces the energy of the depositing atoms must be reduced as much as possible. However, the optimal energy of the metal atoms depends on their position in the stack. Whereas metal atoms below critical interfaces are preferentially deposited at high energy to ensure good surface mobility and smoothness, metal atoms above the critical interfaces are better deposited at lower energy to avoid interface mixing. In the case of IBD, we also have to avoid mixing induced by energetic reflected neutrals which also impinge the sample. This is usually done by the use of a gas which is heavier than the metal atoms, e.g., Xe, so that the reflected neutrals are rapidly slowed down by the heavy Xe ions.
4.6.4. Hard bias magnetic thin films Permanent magnet hard bias films have also found a wide application in giantmagnetoresistive reading heads (Hedge et al., 1999). The general requirements for a hard magnetic material are an adequate coercivity and a large biasing field at small thickness, i.e., high M^ct, where Mr is the remnant magnetisation and t is the film thickness. In magnetoresistive head applications, CoCrPt films with a Cr underlayer deposited by IBD is one of the most frequently employed permanent magnet structures. In fact, it has been found that IBD allows the manipulation of the structural and magnetic properties of Cr/CoCrPt films by properly varying the ion beam energy and deposition angle to achieve a good hard biasing effect. The use of IBD at glancing angles with respect to the substrate plane has been reported (Hedge et al., 1999) to produce Cr/CoCrPt bilayers with low resistivity, high coercive fields Hc and high Mr*.t. The high quality seems to be due to an improvement of the in-plane c-axis texture and the film crystallinity due to enhanced adatom mobility at low deposition angles (cf., section 3.9.6). The possibility of further improvement by the use of an ion assisting source during deposition has also been studied by Hedge et al. (1999), resulting in a coercivity increase as the energy of the assisting ions increases whereas the product Mrxt remains unchanged. The coecivity increase of the ion assisted CoCrPt layers, is believed to be due to the enhancement of the in-plane texture. Coercivities in the range 1400-2100 Oe have been reported for Cr/CoCrPt films (5 nm/25 nm) (Hedge
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Low ENERGY ION ASSISTED FILM GROWTH
et al., 1999). The coercivity increases as the angle of incidence of the deposition beam increases and as the assist energy increases. These trends are attributed to an increase of the (1010) in plane c-axis texture. Another example of the effect of low energy ion beam assisted deposition on the microstructure and magnetic properties of Co^Ptj alloy films has been reported by Sharma et al. (1999). Interestingly, these authors have shown that the use of an ion beam to assist the deposition modifies significantly the magnetic properties of the Co-rich films but has no noticeable effect on the properties of films with x > 0.5. Figure 4.11 shows the variation of coercivity as a function of the ion beam energy for three different compositions (x = 0.2, 0.5 and 0.75). For films with x = 0.2 and 0.25 (not shown) the coercivity shows a maximum of 2 kOe at an ion energy of 250 eV. By contrast the coercivity of films with x = 0.5 and 0.75 remains small and unaffected by the ion bombardment.
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200
300
400
500
Ion beam energy (eV) Figure 4.11. Variation of the coercivity as a function of the ion beam energy for Coi-xPtx for a) x=0.2, b) x=0.5 and c) x=0.7. Reproduced from Sharma et al. (1999), with permission.
The in-plane and perpendicular hysteresis loops are shown in Fig. 4.12. The figure indicates that the magnetic easy axis of the films changes from in plane (non-assisted sample) to out of plane. The origin of those effects is suggested by the analysis of the microstructure of the films. The authors suggest that for x < 0.25 ion bombardment can promote a FCCoHCP transformation, which cannot be induced at higher Pt compositions. In addition, the use of IBAD with increasing energy
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APPLICATIONS OF IAD PROCESSING
induces film growth along the closed packed direction and a gradual orientation of the c-axis out of the film plane.
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Prengel, E.G., Jindal, P.C., Wendt, K.H., Santhanam, A.T., Hedge, P.L., Penich, R.M., Surf. Coat. Technol. 139 (2001) 25. Qin, W., Howson, R.P., Akizuki, M., Matsuo, J., Takaoka, G., Yamada, I., Mat. Chem. Phys. 54 (1998) 258. Renevier, N.M., Lobiondo, N., Fox, V.C., Teer, D.G., Hampshire, J., Surf. Coat. Technol. 123 (2000) 84. Sansom, D., Alonso, F., Ugarte, J.J., Zapirain, F., Oriate, J.I., Surf. Coat. Technol. 84 (1996) 480. Scaglione, S., Flori, D., Soymie, I., Piegari, A., Thin Solid Films 214 (1992) 188. Schmeusser, S., Rupp, G., Hubert, A., J. Magn. Magn. Mat. 166 (1997) 267. Sharma, N., Casey, S.M., Jones, G.A., Grundy, P.J., J. Magn. Magn. Mat. 193 (1999)93. Souche, D., Brunet-Bruneau, A., Fisson, S., Nguyen Van, V., Vuye, G., Abeles, F. and Rivory, J., Thin Solid Films 313-314 (1998) 676. Stippich, F., Vera, E., Wolf, G.K., Berg, G., Friedrich, Chr. Surf. Coat. Technol. 103-104 (1998) 29. Su, Y.L., Yao, S.H., Leu, Z.L., Wei, C.S., Wu, C.T., Wear 213 (1997) 165. Takashashi, H., Appl. Opt. 34 (1995) 667. Tsai, R.-Y., Shiau, S.-C, Lee, C.-H., Ho, F.C., Hua, M.-Y., Opt. Engn. 36 (1997) 3433. Tsai, R.-Y., Shiau, S.-C., Lee, C.-H., Opt. Engn. 37 (1998) 1475. Tompkins, H.G., in A user's guide to ellipsometry (Academic Press, N.Y. 1993) Appendix B. Vera, E., Wolf, G.K., Nucllnstr. andMeth B 148 (1999) 917. Wolf, G.K., J. Vac. Sci. Technol. A 10 (1992) 1757. Wolf, G.K., Preiss, G., Guzman, L., Surf. Coat. Technol. 128-129 (2000) 28. Wolf, G.K., Preiss, G., Miinz, R., Guzman, L., Nucl. Instr. and Meth. B 175-177 (2001) 756. Wang, X., Zhang, F., Li, C , Zheng, Z., Wang, X., Liu, X., Chen, A., Jiang, Z., Surf. Coat. Technol. 128-129 (2000) 36. Yang, Z.W., Han, S.H., Yang, T.L., Ye, L., Zhang, D.H., Ma, H.L., Cheng, C.F., Thin Solid Films 366 (2000) 4.
CHAPTER 5 DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
The interest in developing IAD methods during the last 30 years has been due to their ability to improve the quality of the films via non-thermal mechanisms (cf., Chapter 3). Many experiments and theoretical efforts were also dedicated to the synthesis of new metastable materials, which cannot be made by conventional techniques without the assistance of energetic particles. Since the pioneering work by Aisenberg and Chabot (1971) reporting insulating and hard carbon coatings obtained directly from energetic carbon ions, and the study by Spencer et al. (1976) revealing the presence of microcrystalline diamonds embedded in an amorphous carbon matrix, as well as that by Weissmantel (1981) on the synthesis of cubic boron nitride, the synthesis of both diamond-like carbon (DLC) and cubic boron nitride (c-BN) films has constituted a great challenge for materials scientists. In general, the synthesis of diamond and c-BN require high pressure and high temperature, but after the experiment done by Aisenberg and Chabot (1971), their synthesis also appeared to be achievable by IAD thin film deposition methods under meta-stable conditions. Nevertheless, whereas pure well crystallized diamond films (Grill and Meyerson, 1994; Lee et al., 1999) are usually very successfully obtained by using CVD methods, there is not yet pure chemical methods to deposit c-BN and DLC films. On the contrary, these materials have been synthesized as thin films only with ion-assisted deposition techniques. Several review papers and books cover different aspects of DLC and c-BN deposition methods, structure, diagnostic, properties and applications (e.g., Robertson, 1986; Mishima, 1990; Catherine, 1991; Robertson, 1991; Voevodin and Donley, 1996; Mirkarimi et al., 1997; Grill, 1999) therefore we will mainly consider those aspects and properties which are closely related with the ion bombardment. This chapter focuses on the synthesis of two meta-stable materials, namely diamond-like carbon and cubic boron nitride, for which ion bombardment is a requirement of their synthesis. These two materials are paradigmatic in the sense that direct and quantitative relationships have been established between ion bombardment, their synthesis and their properties. 216
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217
5.1. Diamond-like carbon Diamond is the hardest material known. Its crystalline structure is a zincblende lattice with a four fold sp3 covalent bond structure which gives diamond special mechanical, optical and chemical properties. The hardness, molar density, thermal conductivity, sound velocity and elastic module of diamond are the highest of all known materials, while its compressibility is the lowest. Diamond is transparent over a wide spectral range from IR to UV region (Lee, 1999). All these highly desirable properties have been the motor for an intensive effort to synthesise both diamond and diamond-like materials. Nevertheless, we will focus here on amorphous DLC films, because, as mentioned above, the synthesis of crystalline diamond films is mainly performed by CVD methods and ion bombardment is, in fact, not a requirement for their synthesis. Diamond like carbon (DLC), is the common label assigned to a variety of meta-stable amorphous carbon (a-C) and hydrogenated amorphous carbon (a-C:H) films wim properties such as density, hardness, electrical resistivity and IR transparency, which tend to approach those of diamond. At an atomic level the amorphous DLC films are constituted by a random network of trigonal (sp2) and tetragonal (sp3) and sometimes sp1 coordinated carbon atoms, but in any case with a non-negligible fraction of sp3 bonded carbon. The hydrogen content can reach up to 60 at%. The strong dependence of their structure and physical properties on both the H content and the ratio of sp2 (graphite-like) to sp3 (diamond-like) bonds is in fact, the basis of the well known ternary phase diagram proposed by Robertson (1991) for amorphous carbon materials that has been depicted in Fig. 5.1. The corners of the diagram represent graphite (100% sp2), diamond (100% sp3) and hydrogen. The localisation of a specific diamond-like material on this diagram is determined in terms of the sp2, sp3 and H concentrations which, on the other hand, depend on the deposition technique and the corresponding deposition parameters as will be discussed later. The sp hybridisation is the usual one in graphitic or thermally evaporated carbon films (e-C) with a 90-95 % sp2 bonded carbon. Sputtered a-C is very similar to evaporated e-C in its structure and electronic properties, although it has a higher proportion of sp3 sites (cf., Fig. 5.1). The region labelled as tetrahedral amorphous carbon (ta-C) corresponds to very hard films consisting of mainly sp3-bonded carbon and a density around 3 gem"3, which is along the sp2-sp3 side of the triangle.
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Figure 5.1. Ternary phase diagram of various forms of diamond-like carbon. Reproduced from Robertson (1991) with permission.
Inside the triangle we found two more regions of interest. One is the region labelled as a-C:H corresponding to the hydrogenated amorphous carbon films and characterized by a high concentration of hydrogen and a content of C-C sp3 bonds below 35%. These a-C:H films are softer, less dense and more polymeric-like than the non-hydrogenated a-C films. In fact, it has been observed that the hardness decreases as the hydrogen, content increases, even though the proportion of sp3 sites (e.g., C-H) increases due to an increase in the soft polymeric component. These films are usually prepared by plasma assisted chemical vapour deposition (PACVD) of hydrocarbon precursor gases. In fact, the low content of sp3 bonded carbon could be caused by the relatively low ionisation level of the plasmas commonly used in these methods (Angus and Hayman, 1988). By contrast, hydrogenated a-G:H films with a sp3~bonded carbon of up to 70% and H content of 30% have been obtained by highly ionised mono-energetic plasma beams (Weiler et al., 1996). This material is named hydrogenated tetrahedral amorphous carbon and the corresponding region of existence in the phase diagram has been labelled as ta-C:H. It is seen to lie well above the boundary with the polymers. In the region near the H corner, for hydrogen concentrations above 50 and 60 at% respectively in the sp2-H and sp3-H lines, we get the region of molecules
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which has been labelled as "no film" and HC polymers, where no continuous stable films are possible. Robertson (1986, 1991) has also pointed out that the degree of clustering of the sp2 phase in the film is an important parameter that has not been considered in the phase diagram. In fact, amorphous carbon films with a similar sp3 bonded carbon and H content show different optical, electronic and mechanical properties depending on the size of the clusters of the sp2 phase. According to the phase diagram of Fig. 5.1, it is clear that the higher the content of the sp3-bonded carbon in the deposited DLC films, the more diamond-like the resulting material and their properties will be. This is the reason why there has been a special interest in the search for methods and technologies capable of promoting the sp3 coordination and the diamond-like properties of the films. Both the hydrogenated and non-hydrogenated forms of a-C and ta-C are meta-stable materials that require ion bombardment of the growing films for their synthesis. However, whereas most of the a-C:H films are prepared by plasma assisted chemical vapour deposition of hydrocarbons, the other amorphous carbon forms with much higher content of sp3 C-C bonds (i.e., ta-C and ta-C:H) are usually deposited by direct ion beam or ion assisted PVD methods. These methods provide an effective manner of tailoring the properties of these materials by changing the ratio of sp3 to sp2-bonded carbon.
5.2. Characterisation methods and related properties Considering that the DLC are amorphous materials where the sp3 and sp2 carbon bonded fractions and the hydrogen content are key structural parameters that determine the type and properties of these thin films (cf., Figure 5.1), most of the characterisation methods have focused on the determination of the respective concentrations (i.e., sp3, sp2 and H). This would permit to localise the deposited material in the phase diagram and to evaluate the quality of the diamond-like characteristics of the film. Due to the diversity of closely related materials that can be synthesised, their characterisation is not trivial and requires the complementary use of different techniques to get unambiguous information. In the following, we make a brief presentation of some of the most common characterisation techniques which have proven to give valuable
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information about the structural and bonding characteristics of these materials; namely, structure, sp3 content and density. In the next sections 5.3 and 5.4 we will discuss the deposition technologies and parameters and their influence on the sp3 fraction of the film and on the macroscopic properties (e.g., stress, mechanical, optical, etc.) of the synthesised materials.
5.2.1. Hydrogen concentration The total hydrogen content can be directly determined by forward recoil elastic scattering or by nuclear reaction analysis (NRA) using the 'HC'^, a)12C nuclear reaction. Elastic recoil also provides information about the hydrogen in-depth distribution (Feldman and Mayer, 1986).
5.2.2. Atomic structure (electron and neutron scattering) Electron and neutron scattering are commonly used to determine the bond length, bond angle and coordination numbers in amorphous materials. However, due to the low sensitivity and less availability of neutrons, electron scattering is the routine technique to determine the structure of a-C films. Due to the amorphous character of the a-C films, the structure is of shortrange nature and is described by means of a radial distribution function G(r) which gives a measure of the most probable distances between atoms. Experimentally, it is determined from electron diffraction experiments after measuring the electron scattering intensity to obtain the interference function J(k) (Robertson, 1986). The radial distribution function G(r) for ta-C (90% sp ) has been reported by McKenzie et al. (1991). Its analysis evidences the tetrahedral nature of the material as well as many characteristics that compare rather well with crystalline diamond (100% sp3), i.e., a mean nearest neighbour distance of approximately 0.153 nm (for diamond 0.154), a bond angle of about 110° (109.5° for diamond) and a mean coordination number greater than 3.8 (4 for diamond). The interference function J(k) derived from electron diffraction experiments in ta-C:H films has been reported by Frauenheim et al. (1994). These authors developed an atomic model of the structure which fits rather well with reported structural data and diffraction experiments. Furthermore, these authors showed that the values of the bond length
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and bond angle of the ta-C:H films vary continuously with the sp3 and hydrogen content.
5.2.3. sp3/sp2 bonding fraction (Raman, NMR, EELS/XAS) Raman is the most popular non-destructive technique for the characterization of amorphous (DLC) and crystalline diamond films. In principle, Raman should allow the identification of both the sp3 and sp2 sites in DLC films. However, the common use of a visible laser excitation at 418 nm or 514 nm leads to spectra dominated by the sp2 contribution, that overshadows the sp3 contribution even for ta-C films. This is so because the Raman cross section for sp2 is much larger (50-200 times) than for sp3 sites. By contrast, the use of UV-Raman with a laser excitation at 244 nm (i.e., 5.1 eV) is capable of exciting and probing both the sp2 and sp3 sites with comparable sensitivities. The difference is clearly shown in Figure 5.2 where typical UV and VIS Raman spectra of the same ta-C and ta-C:H films have been depicted together for comparison.
500
700
MM) 1100 1300 1500 170(1 1900 HIM) 1200 1300 1400 1500 1600 1700 ISM)
Raman Shift (cm*1) Figure 5.2. Raman spectra for ta-C and ta-C.H films: a) in the UV range (244 nm), b) in the visible range (488 nm). Reproduced with permission from Adamopoulos et al. (1999).
Figure 5.2 a) shows the UV-Raman spectra (excitation source at 244 nm) for ta-C (80% sp3) and ta-C:H (60% sp3) films reported by Adamapoulos et al. (1999). In the case of the ta-C film, the spectrum exhibits a peak near 1100 cm"1 in addition to the so called G-peak at 1650 cm"1. The peak at 1100 cm"1 results from the stretching of sp3 C-C bonds and it is not observed in Vis-Raman. The peak at
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1650 cm"1 is due to the stretching vibrations of sp2 bonded carbon. For the ta-C:H film, the UV-spectrum also consists of two features, the G-peak at 1600 cm"1 associated with sp2 bonded carbon and a broad peak near 1250 cm"1, which is due to the C-C stretching mode of both hydrogenated and non-hydrogenated sp3 sites. The sp3 fraction of the film is then derived from the ratio of the two peak intensities (Gilkes et al., 1997; Adamopoulos et al., 1999). The Vis-Raman spectra for the same films obtained with an excitation of 488 cm"1 are shown in Figure 5.2 b). In this case the spectra are clearly dominated by a broad and intense G-peak at around 1550 cm"1 associated with the vibrations of the sp2-bonded carbon, whereas the expected diamond line (sp3) at 1330 cm"1 is missing even though the sp3 content in both films is higher than 60%. The spectrum corresponding to the ta-C:H film shows a shoulder at around 1350 cm"1 due to the presence of ordered clusters of sp2 bonded carbon (D-peak) (Gilkes et al., 1997; Adamopoulos et al., 1999). Ferrari and Robertson (2000) have performed a detailed analysis of the VisRaman spectra of disordered graphite, amorphous carbon and diamond-like films. This analysis has enabled a classification of all the available Vis-Raman spectra in the frame of a three-stage model of disorder, which provides a way of identifying the factors that control the position, intensity and width of the G and D peaks in the different Vis-Raman spectra. Due to the fact that in as-deposited ta-C and a-C:H films the sp3/sp2 ratio is directly related to the sp2 configuration, the sp3 concentration can be derived from the shift of the G-peak and the ID/IG intensity ratio. However, as the sp2 configuration can be changed independently of the sp3 content by post-deposition treatments and high deposition temperatures, the above analysis and relationship cannot be extended to treated films. Fourier transform infrared spectroscopy (FTIR) is widely used to characterize C-H vibrational modes in polymers and hydrogenated carbon films, but it only probes unambiguously sites bonded to hydrogen. IR spectra of hydrogenated carbon films display a broad absorption band at 2900 cm'1, which although resulting from the superposition of different C-Hn configurations, should in principle enable an evaluation of the different contributions by the corresponding spectral analysis. However since Grill and Patel (1992) showed that neither the sp2/sp3 ratio nor the absolute hydrogen content of DLC films can be estimated from the analysis of FTIR spectra, the use of this technique is limited to examining qualitatively relative changes in a series of films (Paterson, 1996; Weiler et al., 1996).
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Direct evidence of sp2 and sp3 site fractions can also be obtained by 13C nuclear magnetic resonance (NMR) spectroscopy (Golzan et al., 1995; Donnet et al., 1999). The NMR spectroscopy provides well separated peaks for the two carbon hybridisations, sp2 and sp3, and therefore has the ability to determine the ratio of these two types of carbon atoms. However, the use of this technique requires very large or 13C enriched samples and is therefore difficult to perform for most of the deposited films. High energy electron loss spectroscopy (EELS) is one of the most popular techniques to estimate the sp2 fraction of DLC films, since unambiguous proof of the sp sites can be obtained from the analysis of the carbon K-absorption edge. Characteristic EELS spectra of several ta-C films deposited by filtered cathodic vacuum arc (FCVAD) at different bias voltages as reported by Waidmann et al., (2000) have been reproduced in Fig. 5.3. Reference spectra for graphite and crystalline diamond have been included in the inset for comparison.
Energy (eV) Figure 5.3. EELS spectra of different a-C films grown by FCVAD at different bias voltages. Spectra for HOPG and crystalline diamond are included for comparison. Reproduced from Waidmann et al. (2000) with permission.
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The lines labelled from 1 to 3 correspond to the centre of the Is —• n spectroscopic transition in graphite and to the onset of the Is —• o* transitions in diamond and graphite respectively (as indicated in the inset of Fig. 5.3). The spectra are characterised by a peak at around 285 eV due to excitations from the C Is level to empty u*-states associated to sp2 bonded carbon atoms, followed by a broad feature due to excitations from the C Is level to the empty o*-states corresponding to both sp2 and sp3 bonded carbon atoms. In the case of diamond (100% sp3), for which there is no 7t*-states, the first peak is missing and only the C Is —» a transitions are observed. The sp2/sp3 ratio in the films can be estimated from the relative n*/a* integrated areas, assuming that the decrease of the 7t*-peak intensity is caused by the increase of sp3 states. The results are usually compared with a standard sample of disordered graphite (100% sp2). Similar information to that obtained by EELS can also be determined by Xray absorption spectroscopy (XAS) by looking at the carbon K-absorption edge. The similarity between the EELS and XAS spectra is expected as both reproduce the carbon K-absorption edge at about 285 eV. At low energy losses (i.e., < 50 eV) the EELS spectrum is closely related to the energy-loss-function of the material. In general, it is dominated by the plasmon loss peak, whose energy E p is related to the valence electron density (cf., Eq. 5.1 in section 5.2.4). The plasmon energy has also been found to correlate well with the sp3 fraction and even a linear interpolation between the plasmon energy value for graphite (100% sp2), E p = 24 eV, and that for diamond (100% sp3), E p = 33 eV, has often been used to estimate the sp3 fraction in the a-C film. High plasmon energies around 29.5-30.5 eV indicate films with a high sp3 fraction (i.e., ta-C).
5.2.4. Density A property closely related with the configuration of DLC films is their density. Diverse methods have been reported in the literature for its determination. These methods include flotation, weight gain, a combination of RBS and profilometry, ultrasonic surface waves, plasmon energies measured by low energy EELS and Xray reflectometry (XRR).
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According to Ferrari et al. (2000), XRR is the best method to estimate the mass density whenever applicable. XRR measures the X-ray intensity reflected in the specular direction of the incident X-rays as a function of the grazing-incidence angle (cf., section 3.3.6). The measured XRR curves provide information on the density, roughness and thickness of the layers after their analysis and simulation by the appropriate methods (Ferrari et al., 2000). The most popular method of estimating the density of DLC films is that derived from the measurement of the plasmon energy E p by EELS. Although the EELS spectroscopy can be performed in both the reflection and transmission mode, one has to be conscious that the information obtained by EELS in the reflection mode comes mainly from the surface instead of the bulk of the material, and therefore can be affected by the presence of modified surface layers. The low-energy-loss spectrum in EELS is dominated by the plasmon loss peak, whose energy is related to the local valence electron density by:
p
«
Lm£o.
where n is the valence electron density, m the effective mass of the valence electrons, and e„ the permitivity of the free space. Assuming that carbon contributes to the band with four valence electrons and hydrogen contributes with one, the plasmon energy is a measure of the valence electron density n, and therefore an indirect measurement of the atomic density of the material (Ferrari et al., 2000; Libassi et al., 2000). A detailed comparison of densities derived from XRR and plasmon energies has been performed by Ferrari et al. (2000), who have evidenced the existence of a good agreement between both values if the effective mass m of the electrons in carbon films is assumed to be rn = 0.87 me, where me is the free-electron mass. Table 5.1 summarises some densities reported for ta-C films deposited by FCVA at several bias voltages (Ferrari et al., 2000) and ta-C:H films obtained by plasma beam sources (e.g., PBS and ECWR) (Weiler et al., 1995) as well as density data reported for a-C films obtained by DIBS by Lacerda et al. (2000) and
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magnetron sputtering combined with ion plating by Schwan et al. (1996). The sp3 content of the films is also included for a better comparison between different films. Table 5.1. Density and sp3 content for different series of DLC films.
Method, deposition parameters FCVA, single-bend
density g/cm3
sp3 content (%)
-320 V -290 V -200 V -100 V -80 V -35V -10V +10 V ECWR (25% H) (30% H) (40% H) IB AD, Ar + 0eV (I/A-
2.7 2.86 3.03 2.9 3.24 2.8 2.6 2.71
72 76 81 80 87 73 67 73
2.39 2.13 1.6
70 70 65
1.75 2.25 2.25 2.2
<10 35 25 25
3.2 2.5 2.1 1.72
85 50 30 <20
nat. diamond
3.52
100
graphite
2.27
0
sample Ferrari et al. (2000) Waidmann et al. (2000) ta-C
Weileretal., (1996) ta-C:H
Lacerda et al.(2000) a-C
Schwan et al. (1996) a-C
a-C
It is interesting to note in Table 5.1 that the density of the films varies between that for sputtered or evaporated carbon (i.e., around 1.7 g/cm3) and the one
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for diamond (i.e., 3.52 g/cm3), depending not only on the sp3 fraction, but also on the hydrogen concentration. The density of the ta-C films grown under optimised conditions with the highest sp3 content (i.e., >80%) approaches the value of natural diamond. In this regard, Ferrari et al. (2000) have confirmed a near-linear dependence between the density and the sp3 fraction for both ta-C and ta-C:H films. In the case of DLC films obtained by IBAD methods, the low sp3 content and the high density values suggest that these materials are closer to dense graphite-like films than to diamond-like ones, except for those films deposited at very high I/A ratios, where the sp3 fraction is enhanced significantly. a-C:H films obtained by sputtering or PECVD methods show densities around 1.7 - 2.0 g/cm3.
5.2.5. Cross sectional structure and in-depth composition (TEM, EELS) Cross sectional Transmission Electron Microscopy (TEM) in combination with spatially resolved EELS can be used to obtain bright field images and direct information on the composition and bonding of DLC films. In fact, cross sectional TEM images and in-depth analysis by EELS of ta-C films have been published by Davis et al. (1995 and 1998). These authors have demonstrated that the films are in fact not homogeneous but show three distinct layers along the depth of the film. In general, and regardless of the deposition method it is always found low density layers at the top surface and at the film/substrate interface and the high density layer in the middle (Davis et al., 1995, 1998). Furthermore, the thickness of the surface and interface layers have been found to increase with increasing deposition energy. Figure 5.4 reproduces the in-depth distribution of the area densities of carbon and silicon atoms and the fraction of the sp3 bonded carbon obtained from spatially resolved EELS line profiles along the cross section of a ta-C film grown on Si(001). The film was grown using a filtered cathodic arc system with a bias of -320 V and shows oxygen and calcium as impurities. The in-depth profile has been published by Davis et al. (1995 and 1998). From Fig. 5.4 we can distinguish on the left region an interfacial mixing layer of approximately 5 nm at the Si/ta-C interface. It is induced by the impinging carbon ions during the first stages of the DLC growth on Si. Carbon in this layer is sp bonded. Figure 5.4 shows also rather clearly the presence of a surface layer of about 1.3 nm, which is a layer of sp2-bonded carbon. This surface layer is characterised by the rapid decrease of the sp3 bonded carbon, it is indicated by an
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arrow and is a clear evidence that the transformation from sp2 to sp3 bonding is occurring below the surface at the implantation depth of the incident carbon ions during film growth. The bulk of the film is 100% constituted of carbon atoms, from which around 90% are sp3-bonded.
••«• - S i l i c o n —o— Calcium -•— sp3 f r a c t i o n
• * — Carbon —-0— Oxygen
« 20
100%
• * - • - • • « • • • W»7
1
15 -
d ° H
«
10
O •H •P
U
ao u
Pi I inl.il.lii..l.*t]ft'»»?
10 15 Z Position
20 (nm)
25
0% 30
Figure 5.4. In-depth distribution of the sp fraction and carbon, silicon, oxygen and calcium along the cross section of a ta-C film deposited by FCVAD at -320 V. Reproduced with permission from Davis et al. (1998).
5.3. DLC deposition methods In principle, the deposition of DLC films requires a source of energetic carbon species. The carbon source can be an ionised gas containing carbon species or a graphite target which is thermally evaporated, ion sputtered or laser ablated. The required energy is then supplied by direct electrostatic acceleration of the primary beam or by indirect momentum transfer through collisions with secondary energetic species.
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229
The energy of the carbon species plays an essential role in the formation of the sp3-bonded carbon. Its importance has been quite clear since the beginning. Films containing predominantly sp2 bonds are commonly obtained by thermal evaporation or sputtering of graphite, whereby the carbon species have energies between 0.1 eV and 10 eV. Conversely, the deposition of tetrahedral amorphous carbon films (i.e., ta-C) with a high concentration of sp3-bonded carbon atoms, requires carbon species with energies above several tens of eV. Therefore, a variety of deposition techniques which involve energetic species were dedicated and developed to fabricate DLC films with a high sp3 content. Carbon ions with energies between 30 and 1000 eV, which are necessary for the deposition of DLC films, are commonly employed in the form of plasmas and/or ion-beams (cf., chapter 2). From this analysis we exclude chemical vapour deposition methods in combination with some kind of ion assistance which, although they are very successful obtaining a-C:H films, are out of the scope of this book. Ion assisted methods like DIBS, ion plating, ion beam assisted evaporation, etc. (cf., chapter 2) have been widely used to grow a-C films. In all these techniques low energy carbon species condense onto a substrate while a beam of energetic species (100-1000 eV) transfers momentum to the condensed species and supplies the energy required for the formation of sp3 sites. The relevant deposition parameters are then, momentum transfer, substrate temperature, and the ratio of ion flux to the flux of thermal carbon species (i.e., I/A). In many experiments, where the ion mass and ion angle of incidence are maintained constant, the normalised energy per deposited atom can be used as deposition parameter (cf, section 1.8.1). Taking into account the variety of deposition parameters, including the use of different mixtures of assisting ions, as well as the wide range of variation of the parameters and their interdependencies, that characterise many of the IAD methods we can easily understand the reason why such a variety of a-C materials and properties has been reported in the literature. Due to the low efficiency of the indirect energy transfer mechanism during growth, the IBAD methods require a rather high ion to atom arrival ratio (e.g., I/A = 5-10) to achieve a significant (60-80 %) content of sp3 bonded carbon in the growing film. So, some care must be taken to limit the re-sputtering effects to get an acceptable growth rate. Direct ion beam methods like mass selected ion beam deposition (MSIBD) and filtered cathodic vacuum arc deposition (FCVAD) employ energetic C+ ions for the deposition, whereby the C+ ions are accelerated by means of an external voltage, which defines the energy of the ions (cf, section 2.2). These methods are especially
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suitable for the deposition of ta-C films with a high concentration of sp3 bonded carbon. In these methods, basic deposition parameters are the energy of the carbon ions and the substrate temperature, both of them well defined and controllable. At present, the most highly sp3 bonded form of non-hydrogenated a-C (up to 90%), i.e., ta-C, has been deposited by ion beam deposition (IBD) and cathodic vacuum arc deposition (FCVA) using C+ ions with energies around 100 eV and maintaining the substrate at room temperature. Similar ta-C films have been also produced by pulsed laser ablation (PLD) of graphite by appropriate biasing of the substrate (Voevodin et al., 1996). By contrast, a-C films deposited by IBAD show in general a much lower sp3 content due to a less efficient energy transfer. In fact, many of the IBAD a-C films are better described as a highly compressed and dense sp2 network rather than in terms of sp3 bonded carbon (Lacerda et al., 2000). In the case of hydrogenated a-C (a-C:H) films, plasma enhanced CVD methods using hydrocarbon gases seem to be the most successful methods of deposition and, therefore, they will not be considered here. In general, these a-C:H films are characterised by sp bonding fractions ranging between 0.1 and 0.6 and a hydrogen content up to 40% (cf., Figure 5.1). The incorporation of hydrogen stabilizes the sp3 sites (i.e., most of the carbon atoms attached to hydrogen) and therefore constitutes an important parameter which determines many properties of the a-C:H films, (e.g., they become rather polymeric if the hydrogen content is high enough) (cf., Figure 5.1). ta-C:H films with hydrogen contents up to 30 at % have been successfully grown by highly ionised mono-energetic plasma beam sources (PBS) using acetylene (C2H2) or methane (CH4) as the source gas (Weiler et al., 1996; Sattle et al., 1997). In the plasma beam source the average ion energy is well defined and can be controlled by the plasma d.c. self bias. Taking into account that the impinging ionised molecules (e.g., C2H2+) will dissociate on the surface so that the kinetic energy is equally shared by the two carbon atoms and assuming that the momentum of the original molecule is conserved, the energy per incident carbon atom can be easily obtained (Weiler et al., 1996). These ta-C:H films with a sp3 fraction of up to 70% have a similar hydrogen content ranging around 30%. They are also very interesting materials because they keep, to a high extent, the mechanical properties of the previously discussed ta-C films. ta-C:H films with a much lower hydrogen concentration (<15 at. %) have been deposited by using a
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
231
FCVAD system (Kleinsorge et al., 2001) where hydrogen is introduced to the deposition chamber at different partial pressures (<10~3 mbar).
5.4. Influence of the deposition parameters on the sp 3 bonding fraction and related properties Nowadays it is possible to fabricate amorphous carbon films with almost any fraction of sp3 and sp bonds by using a large variety of ion assisted techniques and varying the deposition parameters. The only requirements for a successful deposition of DLC films are a relatively low substrate temperature (i.e., from room temperature up to 100°C) and energetic carbon species, i.e., carbon containing species whose kinetic energy is well above the thermal energy (i.e., >30 eV). There is also an upper limit for the kinetic energy of the deposited carbon atoms (-1000 eV) that is associated with a significant increase in damage (cf., sections 1.6 and 3.6). This section reviews data from different laboratories using different techniques to asses the influence of the most relevant deposition parameters, i.e., ion energy, substrate temperature, angle of incidence of the ions, etc. for both direct ion deposition methods and ion assisted methods.
5.4.1. Influence of the ion energy The energy of the carbon species is a key parameter in DLC deposition. The dependence of the concentration of sp sites on the energy of the carbon ions (i.e., C+) has been well characterized by several groups using direct energy input deposition methods (i.e., MSIBD and FCVAD). Figure 5.5 shows the sp3 bond fraction reported by different groups using MSIBD and FCVAD methods for films deposited at room temperature (cf., section 5.2.3). The data from Lossy et al. (1995), Xu et al. (1996) and Chhowalla et al. (1996) correspond to films deposited by FCVAD, whereas those from Lifshitz et al. (1995), Hakorvita et al. (1995) and Ronning et al. (1997), correspond to films grown by MSIBD. The sp3 fraction was estimated, whenever possible, from the reported plasmon energies assuming a linear relationship between the value for graphite, 24 eV (100% sp2) and the one for diamond, 33 eV (100% sp3) or from the intensity of the it - 7t* feature at the carbon K absorption edge (cf., section 5.2.3) (Hofsass et al., 1998).
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80
60
a !
40-
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20-
200
400 600 800 Ion Energy [eV]
1000
Figure 5.5. Concentration of sp3 bonded carbon atoms in a-C films as a function of the C + ion energy. Adapted from Hofsass et al. (1998).The data depicted are from, Lifshitz et al. (1995), Hakorvita et al. (1995), Lossy et al. (1995), Chhowalla et al. (1996), Xu et al. (1996) and Ronning et al. (1997).
Figure 5.5 shows a general good correspondence in the overall shape of the dependence of the concentration of the sp3 bonded carbon atoms on the energy of the C+ ions, as determined by several independent groups. We can observe a sharp increase up to an 80-85% of sp3 content for energies between 30 and 100 eV, which indicates rather clearly the existence of a threshold energy of several tens of eV for the formation of sp3 bonded carbon. This threshold is followed by a broad maximum at energies between 100 eV and 300 eV and a smooth decrease at energies above 300-600 eV. This decrease suggests that too energetic species (>600 eV) end up in the deposition of increasingly higher fractions of sp2 bonded carbon. An sp3 fraction of 20-30 % has been reported for films deposited at 10 keV (Grossman et al., 1996). The results published by McKenzie et al. (1991), Fallon et al. (1993) and Pharr et al. (1996), corresponding to films deposited by FCVAD have not been included in Figure 5.5 because although they fit rather well with the general behaviour at threshold, they show a more rapid decrease of the sp3 fraction at energies above 200 eV (cf., Figure 5.6 a). Apparently, this behaviour is due to a significant heating of the substrate caused by the much higher ion current densities (e.g., several mA/cm2) commonly employed in FCVAD as compared with MSIBD (cf., section 5.3.2 for the effect of the temperature).
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
233
The consideration of other data (not included) from films deposited by other IAD methods, where the energy to the carbon species is supplied indirectly (e.g., DIBS, IB AD, etc.), results in a much poorer definition of the threshold energy in the low energy region, mainly due to the usually broad spread in energy of the assisting ions, which include very low energy species, and to the large spatial divergence of the most common low energy ion beam sources (cf., section 2.4). In contrast, the energy (~ 100 eV) of the C+ ions at which the maximum sp3 concentration is obtained appears to be well reproduced by almost all the deposition methods. The sp3 fraction in ta-C:H films deposited by using high plasma density sources (e.g., PBS, ECWR) (Weiler et al., 1996; Sattel et al., 1997) shows a similar dependence on the ion energy to the one discussed above for the deposition of ta-C films, indicating that the growth mechanism is the same. The results reported by Weiler et al. (1996) using C2H2+ ions show that the formation of sp3 bonds and dense ta-C:H films requires an energy threshold of about 57 eV per carbon atom, and reaches an optimum of sp3 bonded carbon (80%) at an energy of 92 eV per C atom. In this case, the energy per carbon atom is estimated by assuming that the energy of the molecule (e.g., C2H2+) is shared by all the atoms. In fact, acetylene dissociates into two carbon ions so that the energy per carbon atom is just 46% of the energy of the impinging C2H2+ ions. The hydrogen concentration of ta-C:H films deposited at room temperature is around 25 % and it is almost independent of the ion energy in contrast to the common a-C:H deposited by PECVD methods (Weiler etal., 1996). In summary, it appears well established that at room temperature there is a threshold energy necessary to promote the formation of sp3 C-C bonds. It is also generally accepted that for too high energies (i.e., > 600 eV) the film is damaged and the formation of sp3 bonds hindered. Therefore, we find a broad energy window (50-600 eV) for the deposition of sp3-rich a-C films. An AFM study performed by Lifshitz et al. (1994) has clearly elucidated the energy effect of the C+ species. The AFM images demonstrate that for C+ energies in the range of 30 - 600eV, which is well known to lead to sp3 rich films (i.e., ta-C), the surface remains atomically smooth as a clear indication that the growth process is internal. This internal growth leads to dense films and smooth surfaces. For higher energies the sp3 fraction decreases but the smoothness of the surface remains, except at very high energies around 20 keV, for which the sp3
234
Low ENERGY ION ASSISTED FILM GROWTH
growth is eliminated and surface roughness increases rapidly. If the energy of the carbon ions is too low, i.e., below the threshold at 30 eV, the formation of sp2 is enhanced and that of sp3 inhibited so that the surface evolves to graphitic-like and rather rough as a clear indication that the carbon atoms remain at the surface and that only surface processes are occurring. It is interesting to note that not only the sp3 concentration of the film but a variety of other properties closely related to the sp3 fraction (e.g., nearest neighbour distances and bond angles, plasmon energies, density, stress, optical gap, etc.) present a similar trend with the energy of the depositing species to that shown by the sp3 concentration of the film (Fallon et al., 1993; Weiler et al., 1996; Chhowalla et al., 1997; Siva et al., 1996; Lacerda et al., 2000). As an example, Figure 5.6 shows the variation of sp3 fraction, plasmon energy and stress for ta-C films deposited by FCVA and the evolution of the sp3 fraction, density and stress of taC:H films deposited by a plasma beam source as a function of the bias voltage and ion energy per C atom. The data correspond to values reported by Fallon et al. (1993) (Figure 5.6 a) and Weiler et al. (1996) (Figure 5.6 b) for the respective type of films. In both cases, we observe that these properties are well correlated with each other and pass through a broad maximum at energies around 100 eV. At energies below 90 eV and above 200-300 eV, there is a rapid deterioration of all those properties towards those associated with graphite-like amorphous carbon (i.e., sp3 fraction < 0.4). Figure 5.6 a) shows data from Fallon et al. (1993) that were not included in Figure 5.5, because they show a more rapid decrease of the sp3 content at energies above 100 eV than the general behaviour depicted in that figure. The stress as well as the optical and mechanical properties of these type of films will be discussed in (cf., sections 5.5 and 5.6).
5.4.2. Influence of the substrate temperature It has been observed that the sp3 fraction of both ta-C and ta-C:H films is strongly dependent on the deposition temperature, changing rather sharply to sp2 above a transition temperature between 70 and 300 °C, depending on the ion energy.
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As pointed out in the previous section, other properties closely related with the sp content show a similar dependence on the substrate temperature. As an example, Fig. 5.7 shows the variation of the sp3 fraction, density and compressive stress as a function of the substrate temperature for ta-C films deposited by FCVA withC ions of 90 eV (Chhowalla et al. 1997). Initially, these properties are found to be independent of the deposition temperature but to fall sharply at a transition temperature of 200 °C. This indicates that the films undergo a transition from ta-C to essentially sp bonded a-C if the substrate temperature is maintained above that transition value. The same trend has been observed for ta-C films deposited by MSIBD (Lifshitz et al., 1994) or pulsed laser ablation (Silva et al, 1996), although the reported transition temperature and sharpness of the transition change depending on the deposition method and the deposition parameters.
236
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The variations in sp bonding fraction in ta-C films as a function of the substrate temperature as reported by Silva et al. (1996), Hirvonen et al. (1997), Lifshitz (1996) and Chhowalla et al. (1997) have been plotted together in Figure 5.8. Figure 5.8 shows that the sp3 content of the film decreases rather steeply if the substrate temperature is above the corresponding transition temperature as measured by different authors under different experimental conditions. This transition temperature has been reported between 70 and 300 °C depending on the ion energy and other deposition parameters (e.g., deposition rate). In Figure 5.8, the range is indicated by two vertical lines.
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
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In the case of ta-C:H films, the results are very similar to those for the nonhydrogenated films. The transition temperature for the transformation from sp3 to sp2 bonding has been reported firstly at 260°C, i.e., higher than that reported for non-hydrogenated films (Sattel et al., 1997). Furthermore, the H content in the film has been observed to decrease firstly at temperatures around 100°C due to the release of loose bonded hydrogen, and more rapidly at a deposition temperature of about 450°C associated to the general loss of hydrogen in the film (Sattel et al., 1997). In any case, the existing experimental results show that when the substrate temperature is sufficiently high, i.e., >100-200°C, the sp3 content of the ta-C films decreases very sharply. This dependence on the deposition temperature is the main reason why ta-C and ta-C:H films are usually deposited at room temperature and the effect of the temperature is very often ignored. In practice, those deposition methods that utilise high current densities at high energies (e.g., FCVAD) and therefore high deposition rates, require some active cooling of the substrate to avoid a rise of the substrate temperature because it can strongly influence the sp3 content of the films. Nevertheless, that strong dependence of the sp3 content on the substrate
238
Low ENERGY ION ASSISTED FILM GROWTH
temperature during deposition does not have to be confused with the high thermal stability of the a-C films. In fact, the sp3 fraction remains unchanged when the a-C and a-C:H films are heated up to 650°C in vacuum after deposition. The effect of the substrate temperature has also been observed by AFM (Lifshitz et al. 1994; Sattel et al. 1997). The well-known effect of an increasing sp2 fraction in the film when the deposition temperature is increased has been associated with the evolution of rough surfaces and the migration of carbon atoms implanted at interstitials towards the surface as the film transforms to be more graphitic-like. As the energy of the C+ ions and the implantation depth increase, the temperature required to induce the out-diffusion and transformation to sp2 films must also increase.
5.4.3. Influence of other deposition parameters Several assessments of the effect of other relevant deposition parameters such as the deposition rate and angle of incidence of the impinging ions have been also reported in the literature (Lifshitz, 1996; Silva et al. 1996; Chhowalla, 2001). In the case of the deposition rate, we can expect that the transition temperature will be higher as the deposition rate increases because the thermally assisted out-diffusion rate of interstitials has to compensate the higher deposition rate. Another effect which can occur is an increase of the substrate temperature during the deposition process and therefore a decrease in the sp3 content of the film. This effect has been claimed previously to explain the differences in concentration of sp3-bonded carbon between films prepared by FCVAD at high deposition rates and that in films prepared with a low deposition rate method like MSIBD, even though the energy of the C+ species was the same. With regard to the effect of the angle of incidence of the assisting ions, it is expected that as the angle of incidence increases the penetration of the bombarding species becomes shallower and the sputtering yield increases significantly. Obviously, both effects can influence indirectly the sp3 content of the film and therefore should be taken into account when comparing films from different laboratories using different methods and/or different deposition systems. In general, the use of bombarding species at glancing angles and high energies leads to films with a smaller sp3 content.
239
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
The influence of the type of substrate and the deposition atmosphere has also been addressed but not really studied and, therefore, remain as issues to be investigated if conclusions are to be reached.
5.5. Stress in DLC films In general, films deposited by IAD methods develop compressive stress whose final value depends on the deposition parameters (i.e., ion bombardment and growth temperature) and the film thickness (cf., section 3.10.5). Stress in thin films leads to shear forces at the interface film/substrate which are proportional to the film thickness. When a critical shear stress is reached crack formation and delaminating of the film may even occur. This general behaviour is also found in DLC films. The intrinsic stress in both ta-C and ta-C:H films has been found to be proportional to their sp3 content showing a strong peak (10-12 GPa) at energies around 100 eV per C atom. Figure 5.9 shows that linear dependence between the stress and the sp3 content for different ta-C:H and ta-C films as reported by Weiler et al. (1996). The figure includes data reported by McKenzie et al. (1991) and Fallon et al. (1993) that follow the same relationship. 1 • A o a
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14
240
Low ENERGY ION ASSISTED FILM GROWTH
The stress in the DLC films evolves during film growth and is strongly dependent on their thickness for both hydrogenated and non-hydrogenated ta-C films. Figure 5.10 shows the evolution of the compressive stress in different ta-C films as reported by Chhowalla (2001). For comparison, stress values of about 2 GPa, (i.e., independent on the thickness) for a-C:H films have been also included. The stress values reported in the literature for a-C:H films range between 0.5 and 3 GPa depending on the deposition technique and parameters. Apparently, the incorporation of hydrogen leads to the formation of C-H bonds, which relieve stress in the film and produce a polymer-like soft material. On the other hand, for the two types of highly tetrahedral carbon films (i.e., ta-C and ta-C:H films) the stress rises up rather steeply when their thickness exceeds 10 nm, and saturates at around 10 GPa for a film thickness between 20 and 100 nm. If the thickness exceeds 100 nm delaminating occurs. During the initial growth of DLC films and due to the ion assistance, a carbide layer is usually formed by ion-mixing at the interface with the substrate (e.g., Fe, Si, Ti, etc.). In fact, this interlayer should give good adhesion to the coatings (cf., section 3.11). However, the high internal compressive stress that develops during the growth of highly tetrahedral films, limits the thickness of the films to much less than 1 um. The reduction of the intrinsic stress to enable the growth of well adhered thicker films is one of the main concerns in several of the fields of applications of these materials. In fact, a considerable number of papers have been published concerning the growth of thick ta-C films through different alternatives to reduce the stress in the film (e.g., Ziegele, et al., 1997; Hirvonen et al., 1997; Friedmann et al., 1997; Chhowalla, 2001). The suggested alternatives include the growth of ta-C/a-C multi-layered structures comprising soft and hard materials by changing the deposition parameters during growth, the incorporation of intermediate layers with Ti or other transition metals, and different post-deposition treatments like annealing or ion irradiation, but a practical and definitive solution has not yet been found (cf., section 5.6).
5.6. Properties and applications of DLC films DLC films represent a group of materials with a broad range of properties depending on the final structure of the film. The properties range from graphite-like to polymer-like and diamond-like, depending only on the deposition techniques and deposition parameters (i.e., energy of the ions, deposition temperature and hydrogen
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
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content). This sensitivity to the processing conditions provides a method of tailoring the properties adapted to specific applications by adjusting the sp2/sp ratio and the hydrogen concentration. 12
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In general, the DLC films may show properties such as high hardness, high wear resistance and low friction coefficients, chemical inertness, IR transparency, high electrical resistivity, field emission properties and low dielectric constant, which enable a wide range of tribological, optical, electronic and biomedical applications (Zellama, 1999, Grill, 1999, Lettington, 1998, Bhushan, 1999). In any case, it is interesting to note that even though ta-C and ta-C:H films present fractions of sp3 bonded carbon (up to 85%) much higher than the a-C:H films (up to 30%) and therefore have properties similar or even superior to those of these films, only these latter appear to have found so far practical applications in industry (Lettington, 1998, Grill, 1999). In fact, the practical applications of ta-C are rather limited and have not yet been proved (Bhushan, 1999, Robertson, 2001). By contrast, a-C:H films appear to have reached maturity in finding practical applications and, therefore, most of the issues discussed below will refer to this type of a-C:H films.
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Low ENERGY ION ASSISTED FILM GROWTH
5.6.7. Mechanical and tribologicalproperties In general the mechanical and tribological properties of DLC coatings are observed to vary strongly. This is not surprising due to the large variety of techniques and conditions used for their deposition. The use of multilayer structures, alloying elements and inter-layers promoters of adhesion results in an even wider range of interesting properties. This makes it difficult to establish clear trends and relationships. Recent reports about a-C:H coatings and their properties can be found in the reviews by Lettington (1998), Grill (1999) and Bhushan (1999). In general a-C:H films have hardness in the range 10-25 GPa and a Young's modulus around 6-10 times larger (Robertson, 1992). The hardness decreases when the hydrogen content increases and the film becomes more polymer-like, even though the amount of sp3 sites may also increase. The reported friction coefficients (Grill, 1997, Donnet and Grill, 1997, Bhushan, 1999) of a-C:H coatings are in the range of 0.007 and 0.4 in vacuum (i.e., <10"4 Pa), and between 0.05 and 1 in ambient air at different relative humidity (20-60%), depending on the deposition method, counterpart material and test environment. Interestingly, the combination of a hard surface with a low friction coefficient and high wear resistance makes these coatings very appropriate for sliding wear protection. Due to its good tribological and wear and corrosion protection properties these films have been used by Gillet® to coat a razor blade introduced in 1998. Nevertheless, it seems that the largest industrial application of a-C:H coatings is as wear and corrosion resistant coatings in magnetic devices (e.g., disks, tapes, read/write heads) (Bunshah, 1999). Another interesting property of the a-C:H films is its biocompatibility. In this respect important applications are being developed for its use in biological environments as orthopaedic prostheses or to protect biological implants against corrosion. a-C:H films deposited on stainless steel and Ti-alloys have been used for components of artificial heart valves and other prostheses, satisfying both mechanical and biological requirements (cf., section 4.3). ta-C and ta-C:H films have extraordinary mechanical properties, such as high hardness and low friction coefficient. The high fraction of sp3 bonded carbon gives these materials mechanical properties which resemble those of natural
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
243
diamond and therefore much higher than those presented by the a-C:H films. Reported values of the hardness for ta-C films are in the range of 40-80 GPa (Schultrich et al., 1998) with Young^s modulus up to 900 GPa. In the case of ta-C:H films, the values reported by Weiler et al. (1996) give a maximum hardness of around 61 GPa and a Young's modulus of 288 GPa for films deposited using a plasma beam source with C2H2 at the optimum ion energy of 90 eV per C atom. Both types of films can therefore be considered as super-hard materials (i.e., hardness > 40 GPa). Unfortunately the high compressive stress of up to 10-13 GPa developed by both types of films limits their thickness to only several hundreds of nm, whereas well adhered films of several um are usually required for most mechanical applications.
5.6.2. Optical and electronic properties Grill (1999) and Lettington (1998) have published extended and recent reviews on the optical and electrical properties of DLC films. Therefore, we only include here a summary of the most relevant properties for completeness. a-C:H films are transparent in the IR, slightly absorbing in the visible and increasingly absorbing as we move into the UV. The index of refraction is found to depend on the preparation conditions and the hydrogen content, so that in principle it can be varied between 1.7 and 2.4 (at a wavelength of 632.8 nm) by changing the deposition conditions. Depending on the deposition system, the reported values of the optical gap (Eg) vary within a wide range even for samples deposited under similar conditions (Grill and Meyerson 1994). Due to its IR transparency a-C:H coatings are being used as antireflective and scratch resistant wear-protective coating for IR optics (wavelength 8-13 um). They have also been used for protection against the scratching of sunglasses lenses made of polycarbonate. The optical properties of ta-C and ta-C:H films deposited by FCVAD and PBS respectively have been studied in detail by Chhowalla et al. (1997) and Weiler et al. (1997). In both types of films the optical gap, the refractive index and the resistivity reach maximum values at an ion energy around 100 eV, in good agreement with a maximum in the sp3 fraction. However, as the optical gap of all forms of DLC are expected to be determined by the gap between the n and the it* states associated with sp2-bonded carbon atoms, we should expect that the gap varies with the sp2 fraction. In fact, this dependence has been observed
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experimentally by several authors. Such a correlation is clearly evidenced by Figure 5.11, where the optical band gap of different DLC films (i.e., a-C:H, ta-C and taC:H) is found to vary almost linearly with the sp2 fraction, indicating that the gap depends primarily on the sp2 sites and very weakly on the hydrogen content.
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In general, DLC films have a band gap which ranges from 1 to 4 eV, but a large number of states in the gap (1021 cm"3) limit rather severely their use and application as amorphous semiconductor (Grill, 1997). The electrical resistivity of DLC films is also generally high, with values ranging between 104 and 108 £2 m, depending on the deposition parameters. Electron field emission of DLC films has been demonstrated experimentally. Both ta-C and ta-C:H have been proposed as low electron affinity materials for the fabrication of large area field emission displays (FEDs). For ta-C films containing a fraction of about 80% of sp bonds, the threshold field for electron emission has been estimated about 107 V/m with the possibility of it being significantly reduced by up to 210 6 V/m by nitrogen doping (0.4 at%). However, a satisfactory explanation of the observed field emission is still lacking (Grill, 1999). Special attention has recently been addressed to the use of DLC and FDLC (fluor-DLC) films as low k materials for inter-connect structures of ULSI circuits to
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
245
improve their performance. By adjusting the deposition conditions, it was shown that DLC films with k values in the range 2.7-3.8 and FDLC with k < 2.8 can be obtained. These materials compared rather well with the value k = 4 of the presently used Si0 2 (Grill, 1999). Obviously, a large effort is being dedicated to optimise properties and to find applications for the DLC materials in general and for ta-C films in particular. However, many such applications have yet to be demonstrated and unfortunately the industrial applications are still rather scarce.
5.7. Cubic Boron nitride films The discovery of the diamond synthesis within a meta-stable regime resulted in an increased effort for the deposition of c-BN under similar conditions. Boron nitride is iso-electronic with elemental carbon and, therefore present similar crystal structures: hexagonal (h-BN), cubic zinc-blende (c-BN) and the wurtzite (w-BN). By contrast with graphite, the h-BN hexagons are arranged on top of each other, while the structure for c- and w-BN completely corresponds to the diamond and lonsdaleite structures of carbon respectively. The c-BN is also known as spharelite or 8-BN. The high density forms of BN correspond to the sp3 bonded structures and include both the cubic phase with a zinc-blende crystal structure (c-BN) and a hexagonal phase with the wurtzite crystal structure (w-BN). The low density phases are sp2 bonded and correspond to the hexagonal (h-BN), rhombohedral (r-BN) and turbostratic phase (t-BN), this latter being a disordered form of h-BN. An amorphous phase of BN is also possible. h-BN, r-BN and t-BN are all sp2 bonded layered compounds, which differ in the stacking arrangement of the hexagonal planes. A detailed structural description of all these phases can be found in Kurdymovetal. (1995).
5.8. Characterisation of c-BN A complete characterisation of the BN films requires, in addition to the determination of the stoichiometry, an unambiguous determination of phases and the corresponding structural characterisation. This requires the utilisation of
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complementary techniques such as RBS/ERDA, NRA, FTIR, electron and X-ray diffraction, TEM, EELS and many other that have been presented in section 5.2 for DLC films. The reader is referred to the review published by Mirkarimi et al. (1997) and the references therein for details on experiments up to 1997.
5.8.1. Stoichiometry The stoichiometry constitutes an important parameter for the growth of c-BN. It has been demonstrated that the composition must be close to B/N ~ 1 for a successful synthesis of the c-BN phase. The most common methods of determining the stoichiometry and homogeneity of the films are RBS and NRA, although Auger and X-ray Photoemission spectroscopies (AES and XPS) in combination with ion beam sputtering have also been used. However, RBS and NRA are a better choice because the surface analytical spectroscopies (e.g., AES and XPS) give only information about the surface composition and, therefore, the results can be affected significantly if the surface has been modified by exposure to air or ion bombardment.
5.8.2. XRD diffraction Due to the crystalline character of the c-BN phase its identification in the BN films can be realised by XRD by comparison with the diffraction pattern of standard cBN. Table 5.2 summarizes the expected diffraction peaks, the spacing and their relative intensities for both the c-BN and h-BN phases. However, the nanocrystalline (5-100nm) and highly defective character as well as the crystallographic disorder of most of the actual c-BN thin films may lead to weak diffraction intensities and the overlap of peaks from the different phases. All of which makes the identification of the c-BN phase in an actual film difficult and sometimes ambiguous. X-ray diffraction experiments of BN films have also been used to study in plane strains of different structural phases, namely t-BN and c-BN, grain sizes and textures.
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5.8.3. Phase identification by FTIR spectroscopy and EELS/XAS In contrast to the problems usually found in the identification of the sp3 bonded carbon in DLC films, the c-BN phase is easily proven by FTIR spectroscopy. In fact, it is the most common method to establish the presence of c-BN. The frequencies of vibration of the infrared active phonon-modes for the cand h-BN phases are well known and allow a rapid screening to determine the local hybridisation of the B-N bonding. For a highly crystalline h-BN phase the characteristic transversal and longitudinal (i.e., TO and LO) modes appear at 780 cm'1 (out of plane B-N-B bond bending) and 1370 cm'1 (in plane B-N stretching mode). On the contrary, the absorption spectrum of a c-BN single crystal shows a peak at 1060 cm'1, which corresponds to the TO phonon mode, whereas the LO component appears at 1310 cm'1. In the case of poorly crystalline BN thin films the FTIR peaks are shifted with respect to the values mentioned above due to the presence of defects and compressive stress. Table 5.2. Diffraction data from the JCPDS-Internafional Centre for c-BN and h-BN.
c-BN [JCPDS 25-1033]
d(A) 2.088 1.808 1.279 1.090 1.044 0.904 0.830
I 100 2 6 3 1 1 3
hkl 111 200 220 311 222 400 331
h-BN [JCPDS 45-0896]
d(A) 3.33 2.17 1.82 1.67 1.32 1.25
I 100 21 18 6 3 7
hkl 002 100 102 004 104 110
Figure 5.12 shows characteristic FTIR spectra for different BN films, prepared by ion assisted deposition under different conditions of ion energy, ion to atom arrival ratio (I/A) and substrate temperature, whereby the characteristic absorption peaks for both the h-BN and c-BN phases can be identified. The relative concentration of both phases will depend on the relative intensities of the corresponding peaks. Neglecting the differences in width and
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absorption coefficients, the c-BN content in the film is commonly derived from the ratio of intensities of the main absorption peaks, making use of the formula: '1080
%(c-BN) = horn
+
-xlOO
A390
.3
I
1 £
600 800 1000 1200 1400 1600 1S00 2000
Wavenumber (cm) Figure 5.12. FTIR spectra of BN films grown by BAD on silicon as a function of a) the ion energy, b) temperature of the substrate and c) the ion to atom arrival ratio. Reproduced from Zeitler et al. (1999) with permission.
(5.3)
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Raman can also be used to determine the phase composition of the films. However, due to the nano-crystalline character of the films and the large number of defects in them, the Raman peaks are usually very broad and of difficult analysis. Evidence of sp2 bonds can be obtained by techniques such as electron energy loss spectroscopy (EELS) or X-ray absorption spectroscopy (XAS), as was shown for the characterisation of DLC films. The near edge fine structure of the two (B and N) K edges can be used to determine the content of sp2-bonded material. Figure 5.13 shows characteristic EELS spectra in the transmission mode for h-BN and c-BN at the boron K-edge. The interpretation is rather similar to that presented in section 5.2.3 for DLC films. For the sp2 bonded h-BN film the spectrum of the boron K-edge consists of a sharp peak at 188 eV corresponding to transitions from the Bis to the anti-bonding JI *states and a broader peak due to excitations to the a* anti-bonding states. By contrast, the spectrum for the sp3 bonded material (c-BN) still shows the broad peak associated with the o anti-bonding states but loses the sharp peak at the edge because there are not n '-states in sp3 bonded materials. A quantitative estimation of the sp2 fraction can be done in terms of the relative intensity of the JI peak in a particular spectrum as compared with its intensity in the spectrum of a pure h-BN (100% sp2) reference sample (cf., section 5.2.3).
-d 9
1
1 170
190
210
230
250
Electron Energy Loss (eV) Figure 5.13. EELS spectra at the boron K-edge from h-BN and c-BN films.
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The large energy difference between the plasmon energy losses in c-BN (-30 eV) and h-BN (-21.5 eV) indicates a large difference of density between these two materials (Feldermann et al., 1999). The actual value for a given film is also used to estimate its density and the c-BN concentration (cf., sections 5.2.3 and 5.2.4).
5.8.4. Microstructure by TEM The microstructure of the BN films can be assessed by transmission electron microscopy (TEM), both using bright/dark field imaging and high resolution observations of the cross sections as well as selected area electron diffraction (SAED). High resolution TEM allows to identify both the sp2 and sp3 bonded nanocrystals of BN films. Regardless of the growth method, the cross sectional TEM pictures of c-BN films reveal a layered structure such as that presented in Fig. 5.14. Figure 5.14 shows the cross section TEM image of a BN film deposited by MSIBD at 350 °C, as reported by Hofsass et al. (1997). The image shows rather clearly the presence of a mixed interface formed by deposited B and N atoms and by Si atoms from the substrate (i.e., Si-B-N) followed by a 10 nm thick layer of a highly textured hexagonal t-BN layer on top of which the nucleation and growth of sp3 nano-crystalline c-BN occurs. The interface layer (t-BN) is textured so that the basal (0002) planes are in the direction perpendicular to the substrate. The TEM image shows that the (111) planes of the c-BN phase which grow on top of the t-BN layer are perpendicular to the substrate or equivalently parallel to the basal (0002) planes of the t-BN interlayer. A detailed observation of this h-BN/c-BN interface also indicates that there is an almost perfect 2:3 lattice matching between the h-BN (0002) and the c-BN (111) planes (cf., inset of Figure 5.14). The formation of the textured turbostratic t-BN layer with the appropriate orientation and inter-planar distance seems to be essential for the nucleation of the c-BN phase (Kester et al., 1994; Medlin et al., 1996; Weissmantel and Reisse, 1999). In addition, it has been found that the surface region of all investigated films is a hexagonal (h-BN) layer with a thickness which depends on the energy of the ions (Park et al., 1997; Hofsass et al., 1997). We should recall here that a similar layered structure including the sp2 surface layer was also observed in the amorphous ta-C and ta-C-H films (cf., section 5.2.5).
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Figure 5.14. TEM micrograph of a BN film deposited by MSffiD at 350 °C and Eton = 500 cV. Reproduced from Hofsiss et al. (1997) with permission.
S3. c-BN deposition methods The synthesis of c-BN under thermodynamic equilibrium requires similar high temperature and pressure conditions as those used to produce synthetic bulk diamond. However, after the successful synthesis of diamond-like carbon films by IAD techniques, Weissmantel (1981) reported the deposition of hard, transparent BN films with a range of properties analogous to those associated with diamondlike carbon films. Ion assisted evaporation was one of the first techniques employed in its synthesis. In fact, Satou and Fujimoto (1983) were pioneers demonstrating that c-BN might be obtained by 30 keV N2+ bombardment of evaporated boron. The resulting films showed the formation of h-BN and some c-BN micro-crystals embedded in a boron rich amorphous phase. Since then, the low pressure synthesis of c-BN thin films has been attempted by many different processes but so far, has only been successful by ion assisted methods, e.g., plasma assisted CVD (PACVB), ion beam assisted evaporation (IBAD), ion assisted pulsed laser deposition (IAPLD), bias magnetron sputtering, mass separated ion beam deposition (MSIBD),
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filtered vacuum arc deposition (FCVAD), etc. (Mirkarimi et al., 1997 and references therein). c-BN films have mostly been grown by IBAD methods using evaporated or sputtered boron atoms and additional ion bombardment, typically with a mixture of Ar+ and N2+ ions (Inagawa et al., 1987; Ikeda et al., 1990; Watanabe et al., 1991; Kester et al., 1994). In the case of IBAD methods, the key deposition parameters controlling the growth and phase evolution in the BN system are the substrate temperature, normalized ion energy or momentum transfer and ratio of ions to condensed boron atoms (cf., section 1.8). As in the case of the deposition of DLC films, there have been several studies on the growth conditions of c-BN using MSIBD of boron and nitrogen ions (Hofsass et al., 1997). Its major advantage over the IBAD techniques is that both nitrogen and boron are directly deposited as singly charged ions of well defined energy (e.g., a few hundred eV) and no noble gas or other species are involved in the deposition. In this case, the deposition parameters are the ion energy, Eion, the B+/N+ ion flux ratio and the substrate temperature.
5.10. Influence of the deposition parameters The success of the synthesis of c-BN films is clearly determined by the bombardment with energetic particles during film growth. Several groups have studied the influence of the different deposition parameters during the growth of the c-BN phase. Iganawa et al. (1987) were pioneers in observing that the growth of cBN was influenced by the ion mass and that it required a certain flux of ions with enough energy. These results were confirmed in successive experiments by different groups using different techniques. Kester and Messier (1992) studied the relationship between the growing phase and the momentum transfer of the assisting ions (assuming a single head on elastic collision, cf, Eq. 1.47) and established the existence of a narrow window of momentum transfer where the growth of c-BN is possible. The effect of the substrate temperature was also studied and the existence of a threshold temperature was confirmed. Another detailed study by Mirkarimi et al. (1994) on the role of the ions in the formation of c-BN films by IAPLD also showed that the formation of c-BN scales rather well with the total ion momentum, depending on the substrate temperature and the B:N arrival ratio. Many other authors also found that the energy threshold for c-BN growth scaled with E1/2. The
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
25 3
measured dependence is schematically shown as full lines in the form of a phase diagram in Figure 5.15 for B/N=l using the momentum transfer and substrate temperature as parameters.
600
I 500 h
/
/ /
/
/ /
/
/Regipnof / / Res6uttesrine'
/
/
> 400 £ 3002 200 h B
|
s o S
MSIBN /' \ nucleation
100
Hexagonal i
100
C-BN regime;
|
200 300 400 substrate temperature [°C]
500
Figure 5.15. Nucleation regime of c-BN as a function of the momentum transferred to the film atoms and substrate temperature as deposited by IBAD (solid lines) or MISBD (dotted lines) methods. Reproduced from Ronning et al. (2000) with permission.
The schema shows as solid lines the nucleation regime of c-BN for films deposited by IBAD. It is shown that for values of the transferred momentum between 200 and 400 (eV amu)"2, i.e., the region labelled as mixed, both the cubic and hexagonal phases are formed if the substrate temperature is maintained above 100°C and the B/N arrival ratio is close to 1. For values below 200 (eV amu)"2 only h-BN is formed and above 400 (eV amu)1'2 there is no net film growth due to resputtering of the deposited material. These conditions for the momentum transfer can be translated in terms of I/A or normalised energy. For a fixed ion mass and energy above threshold there is a critical I/A value above which c-BN formation is initiated, a window of I/A values for which a large percentage of c-BN is obtained, and a point at which I/A yields a re-sputter rate equal or larger than the deposition rate and no net film growth is observed. The growth of c-BN films by IBAD methods is thus a function of the substrate temperature, ion energy and flux ratio of
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ions to boron atoms and overall momentum transfer. The experimental results and the phase diagram show that by ensuring the stoichiometry of the films the growth of c-BN is only observed above a momentum transfer threshold of 200 (eV amu) and above a substrate temperature Ts - 100°C. A similar phase diagram has been proposed by Hofsass et al. (1997) in terms of the ion beam energy and substrate temperature for MSIBD deposited BN films. This has been included for comparison in Figure 5.14 as dotted lines after conversion of the ion energy Eion in momentum transfer. Interestingly, Figure 5.14 shows significant differences between both types of deposition methods. Compared with the IBAD films the films deposited by MSIBD show very sharp energy (momentum transfer) and temperature thresholds as indicated by the two perpendicular dashed lines included in Figure 5.15. Furthermore, the phase diagram for MSIBD films does not include any regime for the growth of mixed h-BN/c-BN phase, but conditions above certain thresholds for which growth of pure c-BN occurs. In the case of MSIBD films, both B + and N* ions are deposited simultaneously and, therefore, there is no re-sputter limit and the growth of c-BN is possible even at high energies. In summary, although the mechanisms involved in the nucleation and growth of c-BN are still not completely understood (cf., section 5.13), the growth of c-BN is not a challenge anymore. Its synthesis is possible by fulfilling three conditions regarding the energy of the bombarding ions, the substrate temperature and the composition. The particles impinging onto the growing film should have an energy above a threshold of about 100 eV (momentum transfer above 200 eV1/2 amu"2); the temperature of the substrate should be maintained above 100-125CC during the deposition and the stoichiometry of the growing film should be close to unity (i.e., N/B = 1). In the case of IBAD methods the momentum transfer shows an upper limit of 400 (eV amu)"2 due to re-sputter effects. Hahn et al. (1997) have nicely shown that although the ion bombardment is a general requirement, it seems specially necessary during the nucleation process and can be reduced significantly during growth. In fact, the above mentioned energy and temperature thresholds have been reconsidered due to new experimental results which indicate that these two parameteres are more closely related to the nucleation than to the growth of c-BN. Litvinov and Clark (1997) have shown that once the nucleation of c-BN has been initiated at energies above the corresponding threshold, its growth can be maintained at lower energies down to 60 eV. Using
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IAPLD, McCarty et al. (1996) also observed that once nucleated, the growth of cBN may take place at temperatures as low as 75°C. Even room temperature growth of c-BN has been proven by Feldermann et al. (1999) if the nucleation was first induced under the established conditions, i.e., substrate temperature above 150°C and Eion above 125 eV, and then the substrate temperature is cooled down to room temperature, while the ion energy is maintained above 125 eV. In any case, lowering the growth temperature below the nucleation threshold reduces the size of the growing c-BN crystallites.
5.11. Stress Irrespective of the deposition technique, c-BN is only formed under low energy ion irradiation, which leads to significant compressive stress in the layers (cf., section 3.10). This stress may cause the film to peel off from the substrate for films thicker than around 100 nm, thus hindering industrial applications. In-situ measurements of stress during growth and post deposition treatments have been performed by Fitz et al. (2000) using the cantilever bending principle and a two laser beam deflection approach. The instantaneous stress is calculated from the curvature and the film thickness data. Figure 5.16 shows the evolution of the instantaneous stress during the deposition of c-BN as a function of the film thickness as reported by Fitz et al. (2000). The film was deposited by IBAD using a mixture of N2+ and Ar+ ions to assist the growth (I/A=2, Eion=600 eV and substrate temperature of 600 °C). Figure 5.16 shows a stress development divided into three different regions that can be associated with the layered structure of the films. Firstly, the compressive stress is observed to increase very slightly up to values around 3-4 GPa during the formation of the t-BN layer at the interface with the substrate. For a thickness above 30 nm, when c-BN starts to nucleate and grow, the stress increases more sharply up to values around 7-10 GPa, depending on the deposition parameters. Finally, after coalescence of the c-BN above 40 nm the stress reaches a constant value characteristic of the c-BN layer. Similar in-situ measurements and results have been reported by several groups, confirming both the general behaviour and the quantitative values (Donner et al., 1998; Klett et al., 1999, 2000, 2001).
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30
40
60
Thickness (nm) Figure 5.16 Instantaneous stress of a growing BN film as a function of the thickness. The nucleation and growth of pure c-BN is indicated by vertical dotted lines. Reproduced from Fitz et al. (2000) with permission.
While growing c-BN is not a challenge anymore, its practical use for commercial purposes is clearly handicapped by such a high compressive stress. Regardless of the method of preparation, the unavoidable ion bombardment causes the build up of such a stress that the achievable film thickness is limited to only a few hundred nanometers. Above that limit, the films tend to delaminate. Deposition methods with reduced stress or stress relaxation methods are, therefore, highly desirable. To achieve this goal, there have been several attempts to deposit buffer layers, growing multilayers, some doping or post-deposition treatments like annealing or irradiation with medium energy ions. Donner et al., (1998) have performed detailed measurements of the in-plane strain of mixed turbostratic/cubic films by X-ray diffraction using synchrotron radiation as well as of its relaxation during thermal annealing. These and other results indicate a massive and irreversible stress relaxation at temperatures above 600°C, without significant changes in the grain size (Mirkarimi et al., 1997; Donner et al., 1998). However, and in spite of all these efforts, the deposition of thick (~1 |jm) films of c-BN has continued to be rarely reported in the literature (Litvinov and Clarke, 1999; Boyen et al., 2000).
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5.12. Properties and applications of c-BN films Cubic boron nitride is a promising material for many applications, e.g., mechanical, electrical, optical, etc. according to its properties. The following sections review some of the properties that confer so a high interest on this material
5.12.1. Tribological properties h-BN is a sp2 bonded layered compound iso-structural to graphite and, therefore, relatively soft. In contrast, the bulk c-BN material show hardness well above 40 GPa and Young's modulus around 800-900 GPa (i.e., c-BN is the second hardest material after diamond). Some advantages with respect to diamond are that c-BN is stable even in air up to 1200K without phase transformation or disintegration and that it does not react with ferrous materials even at 1600 K. This makes c-BN superior to diamond as a wear resistant coating for applications with ferrous materials. Hardness measurements in c-BN thin films are not easy due to the limited thickness of the available films, e.g., just a few hundred nm. In general, nanoindentation is the method most commonly used in the literature. However, as many reports omit the details on the indent depth or the c-BN content in the film, the reported values remain dubious and difficult to compare. McKenzie (1993) has reported a value of 58 GPa by nano-indentation for a 150 nm thick c-BN film at an indent depth of 1 lOnm. For an irradiated t-BN film the measured hardness was 20 GPa at an indent depth of 180 nm. Mirikami et al. (1997) performed also nanoindentation in a 700 nm thick c-BN film at an indent depth of 100 nm, obtaining a hardness of 60 GPa in good agreement with the value reported by McKenzie (1993). Due to its good tribological properties and chemical inertness with ferrous metals even at high temperatures, c-BN could be used in tooling cast irons and high speed steels. In fact, a sintered c-BN cutting tool is available and has been used practically. It has demonstrated its effectiveness in cutting hardened steel even under high cutting speed conditions. However, in addition to its high cost, it shows problems associated with poor ductility and it is difficult to shape adequately for certain applications. There are several reports in the open literature demonstrating
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the effectiveness of c-BN coatings in improving the resistance and tool life (e.g., Jin et al., 2000). However, due to the small thickness of the film (< lum) and the weakness of the adhesion to the substrate the wear resistance is still insufficient.
5.72.2. Optical and electrical properties h-BN is birefringent and transparent in the visible. By contrast, the sp3 bonded c-BN material is an insulator (108 Q m) with a wide band-gap (Eg~6.4 eV) and many potentially applicable properties. These include a refractive index n = 2.1 (at X = 600 nm) and a high thermal conductivity(~1300 W m'C" 1 at 25°C). Moreover, c-BN appears potentially adequate for a broad range of electronic applications in high power and high temperature electronic devices, since it has the widest energy gap among the III-V compound semiconductors and, in principle, could be bom p- and n- doped. c-BN has been proposed for FET transistors for high power microwave applications. The main problem is that most of the c-BN films consist of too small crystallites with a high density of defects, that induce high optical absorption, lower carrier mobility and an excessive electrical conductivity, for most industrial applications (Mohammad, 2002). Nevertheless, promising advances in the epitaxial growth of diamond/c-BN (Pickett, 1988, Yoshikawa et al., 1991) and the potential applications of this hetero-structure constitutes a strong driving force for the development of practical devices. The condition of negative electron affinity (NEA) of c-BN has also been investigated by several authors (Pryor 1996; Loh et al., 1998), however, the data are still scarce and require confirmation.
5.13. Modelling the growth of sp 3 bonded materials (ta-C, ta-C:H and c-BN) With regard to the growth of both DLC and c-BN films, the large experience accumulated during the last 30 years has led to some well established experimental facts which appear to be intrinsic of the DLC and c-BN growth and should therefore be explained by any model or simulation developed to describe the growth mechanism of these materials in the form of thin films.
DIAMOND-LIKE CARBON AND CUBIC-BORON NITRIDE FILMS
• •
•
•
259
The synthesis of both DLC and c-BN require the assistance with energetic particle bombardment. DLC and c-BN only nucleate above well defined threshold values for the ion energy. The sp3 fraction in the respective materials shows a rather characteristic dependence on the ion energy. Both c-BN and DLC films show a thin surface layer of sp2 bonded atoms (cf., sections 5.2.5 and 5.8.4). This surface layer proves that the nucleation and growth of these sp3-bonded materials is occurring in the subsurface region at a depth given by the ion range. DLC and c-BN behave differently regarding the influence of the substrate temperature during growth. c-BN nucleates above a well defined threshold value of the substrate temperature and the crystallinity of the films improves by increasing the temperature of the substrate. By contrast, the deposition of amorphous DLC films at any temperature above 150 °C leads to the formation of sp2 structures.
Throughout these years several theoretical models based on preferential sputtering (Reinke et al., 1994, 1995, 1997), compressive stress (McKenzie, 1993, McKenzie et al., 1993, 1996; Mirkarimi et al., 1994) thermal spikes quenching (Hofsass et al., 1998) and subplantation (Lifshitz et al., 1989, 1990; Boyd et al., 1998; Marton et al., 1998) processes have been proposed to describe the formation of both c-BN and DLC films deposited under ion bombardment. In addition .numerous simulations using Monte-Carlo and molecular dynamics methods have been performed to simulate different processes induced by the ion bombardment (e.g., Uhlmann et al., 1998; Kaukonen and Nieminen, 2000), although a straightforward comparison with experimental data is still difficult due to the complexity of the processes involved and the lack of realistic models. All the models consider three different stages with different time scales which correspond to the basic mechanisms of interaction of ions with kinetic energies in the range 10-104 eV: • •
The impinging ions transfer their energy to the target atoms. This stage lasts 10~ - 10_1 seconds as was discussed in section 1.6.2. A rapid thermalisation stage, in which the atoms participating in the collision cascade lose their excess energy to reach thermal equilibrium with the surrounding atoms. This stage is usually treated in the frame of the thermal spike concept.
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•
A long term relaxation, which includes diffusion processes, chemical reactions, phase transformations and stress relaxation.
The precise mechanisms and detailed events involved in every stage, as well as their role and importance in the whole process, are significantly different for the different models.
5.13.1. The preferential sputtering model The model was developed to explain the growth of c-BN. According to this model (Reinke et al., 1994, 1995, 1997), once some c-BN has nucleated, there are critical sputtering conditions (ion energy and flux) at which h-BN is preferentially sputtered with respect to c-BN. Under such critical conditions of ion assistance the c-BN phase would grow more rapidly or would even be the only growing phase. The nucleation mechanism is considered as an independent process and, therefore, is not included in the model formulation. Initially, the estimation of the critical sputtering conditions showed good agreement with existing experimental results for the energy and I/A ratio, however, as the growth conditions were changed by posterior experiments, especially at low energies, the sputtering model became inconsistent with the experimental results. Furthermore, the sputtering model is unable to explain the presence of a sp2-bonded surface layer, since it clearly predicts that the surface should only consist of c-BN.
5.13.2. The stress models A characteristic feature of both the DLC and c-BN films is the high compressive stress (up to 10 GPa) which evolves during their growth. Therefore, models based on stress induced formation or stress stabilisation of the sp3 bonded phase have been proposed by several authors (McKenzie, 1993; McKenzie et al., 1991, 1993, 1996; Mirkarimi et al., 1994). The main idea is that the compressive stress would be the cause of the formation of these sp3 bonded materials instead of an effect. In fact, the stress model proposed by McKenzie et al. (1991, 1993, 1996) suggested that the stress
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within the material and the high temperature provided by the ion induced spikes could supply the conditions (high P-high T) where c-BN is the stable phase. The compressive stress is considered as due to the accumulation of ion induced defects and atomic penning (cf., section 3.10) and estimated according to the static stress model developed by Davis (1993). This model incorporates a defect production mechanism according to the proposal by Windischmann (1987) and a relaxation process associated with the ion-induced thermal spikes. When applied to the growth of c-BN with (I/A=l) the model predicts (cf., Eq. 3.13) a dependence of the stress on ion energy according to a « E - 1 1 6 , which should decrease as the energy increases due to the enhancement of the relaxation processes (i.e., thermal spike). This behaviour is in clear contradiction with the experimental results which show that increasing the ion energy increases the c-BN content in the film. Later on, Mirkarimi et al. (1994) proposed a more dynamic stress mechanism, which included not only the production of defects but also their annihilation at grain boundaries, so that the residual stress is determined by the corresponding balance of the two processes. Mirkarimi et al. (1994) concluded that the maximum stress should scale with the total momentum of the bombarding ions instead of with the momentum transfer in a single binary collision (cf., Eq. 1.47). In this manner the model accounts for the threshold and window of the total momentum per depositing atom that determine the experimental conditions for the growth of c-BN (cf., section 5.10). The main shortfall of these models is their inability to explain the effect of the substrate temperature on the growth of c-BN. Higher T would allow a more efficient defect relaxation and annihilation and, consequently, less stress in the film. Therefore, if the stress is the controlling factor we would observe a more difficult growth of c-BN as the temperature increases, however, the opposite trend is what is actually observed experimentally.
5.13.3. Models involving a thermal spike mechanism The thermal spike model attributes the formation of DLC and c-BN to the rapid quenching of thermal spikes caused by the ion impact. In this model the thermalisation stage is clearly indicated as the crucial stage. The idea that thermal spikes induced by the energetic ions impinging the film would lead to the
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appropriate conditions of high T and high P for the formation of the sp bonded metastable structures was first suggested by Weissmantel (1981). Assuming high cooling rates (i.e., 1014-1015 Ks"1), the meta-stable structure could be preserved during its thermalisation. The condition of high pressure would be supplied by the large fraction of trapped species which lead to high local stresses. According to existing formulations of thermal spikes, the local temperature in the thermal spike region can reach (depending on the energy of the ions) the melting temperature of the solid. Under these conditions, the meta-stable sp3 bonded structures (DLC and c-BN) would arise due to local recrystallisation processes within the volume where the increase in density and compressive stress has been induced. Using the existing theoretical formulation of thermal spikes by Seitz and Koehler (1954), Hofsass et al. (1998) have developed a cylindrical thermal spike model for ion deposition of diamond like carbon and c-BN films. The cylindrical spike model is described in detail in Hofsass et al. (1998) and the reader is referred there for quantitative evaluations and a detailed discussion. The model treats the energy dissipation in the frame of cylindrical thermal spikes created by individual ion impacts. Hofsass et al. (1998) have applied the model to the formation of DLC and c-BN. For these two materials the model is able to predict quantitatively the optimum ion energy and the ion energy range for which dense sp3 phases are formed. In any case, the crucial question about the thermal spike model is whether the thermal spikes at the energies considered here are actually able to reach melting conditions in high melting ceramics like BN. In fact, this has been questioned by both molecular dynamics simulations and by experiments performed in h-BN using high energy ions (Collazo-Davila et al., 1999). Moreover, the formation of thermal spikes in very small volumes would not explain the observation of relatively large crystallites in these films. In addition, the existence of a threshold temperature of around 100 °C for the formation of c-BN is difficult to explain for a model that involves thermal spikes where the material would be melted (>1000 °C).
5.13.4. Subplantation models Lifshitz et al. (1989, 1990) introduced the subplantation model to describe the formation of sp3-bonded carbon in DLC films by direct ion beam deposition at energies of several tens of eV up to several keV. The subplantation process consists
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in the shallow subsurface implantation of low energy ions with the sufficient energy to penetrate beneath the substrate surface. The subplantation effect has been qualitatively confirmed by both simulations and experiments and appears to be generally accepted. The subplantation model as proposed by Lifshitz et al. (1989, 1990) neglects the role of thermal spikes during the short term thermalisation stage which is not considered in the formulation. The final structure of the film is determined by the collisional and long term relaxation stages and considers the preferential displacement of sp2 bonded atoms with respect to the sp3 bonded species as the mechanism for the growth of the sp3 bonded species. Based on some TRIM simulations, Lifshitz et al. (1989, 1990) conclude that if the preferential displacement of one bond type with respect to the other is the crucial mechanism, the optimum conditions for the formation of the sp3-bonded phase would occur for energies between 50 and 200 eV, in correspondence with the experimental observations. In this case the formation of the graphitic phase at substrate temperatures above 125°C is attributed to the diffusion of carbon interstitials (accumulated in the subsurface region) during the long term relaxation stage. Robertson (1993) made use of the subplantation ideas from Lifshitz but included the thermal spike concept to describe the dependence of the sp3 content on the ion energy. In either case, the conversion sp2 —• sp3 happens at the penetration depth of the energetic species where the local density and stress increase significantly. The existence of a threshold energy is then a consequence of the fact that ions with energies below that threshold are not able to penetrate below the surface and will stick at the surface forming loose sp2-bonded species. According to Robertson (1993) the dependence of the sp3 content on the ion energy results from a balance between densification of the subsurface region, which leads to an increasing sp3 content and relaxation towards a sp2 bonded graphitic phase within the thermal spike volume. The model was initially able to explain the results reported by McKenzie et al. (1991), but, due to its oversimplifications, it had more difficulties in explaining the more detailed experimental results obtained by MSIBD. The most recent proposal of the subplantation model is a semi-quantitative generalization of the initial model (Lifshitz et al., 1989, 1990) developed by Boyd et al. (1998) and Marton et al. (1998). The model proposes analytical equations to estimate the fraction of ions penetrating the surface and the rates of defect
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production and radiation enhanced diffusion as the major contributions to the whole process. Below the threshold energy surface growth and surface effects dominate. However, as the energy is increased above threshold, the ions are able to penetrate below the surface. At much higher energies, damage processes become important and have to be considered. The production of sp3 bonded structures is then determined by a balance between subsurface penetration and densification and damage effects of the assisting ions. The incorporation of the production and radiation enhanced diffusion of defects provides also an explanation of the high energy and temperature effects, which are not accounted for by the thermal spike mechanism. The model was able to account for the sp3 fraction of DLC films deposited by MSIBD as a function of the energy of the ions over a broad energy range between 10 eV and 20 keV. The importance of the radiation enhanced diffusion of defects is then associated with the dramatic decrease of both the density and the sp3 content in DLC films deposited at temperatures above 150 °C. It is likely that the present or eventually new analytical models that could be developed will not improve the current understanding of the growth mechanisms of sp3-bonded materials by ion-assisted methods. Therefore, more efforts should be dedicated to perform better MC and MD simulations that have the potentialities to address complex processes like those involved in ion assisted growth methods (Uhlmann et al., 1997; Kaukonen and Nieminen, 2000; Kohary and Kluger, 2001).
5.14. Related materials (CN„ B-C-N) In 1993, P-C3N4 was proposed as a material harder than diamond (Liu and Cohen 1989, 1990). Its synthesis was attempted worldwide using all the available PVD and CVD techniques, including IAD, IBAD, etc. involving nitrogen ion bombardment of a growing carbon film. An extensive review of all these preparation methods has been published by Muhl and Mendez (1999), so that we refer to it for specific references. However, all those attempts to prepare P-C3N4 clearly failed (Matsumoto et al.; 1999). Nevertheless, based on that prediction, and as a consequence of all those attempts, new amorphous carbon nitride films (i.e., CNX) were developed and also intensively investigated. Carbon nitride films (i.e., CNX) with x ranging from near 0 up to around 0.4 have been synthesized by a large variety of techniques, but mostly by IAD and IBAD methods (Muhl and Mendez, 1999). In general, the films deposited at low
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temperature (i.e., 100°C) are amorphous, but at higher temperature a fullerene-like microstructure has been observed by TEM (Sjostrom et al., 1995; Hayashi et al., 1999). From the paper by Hellgren et al. (1999) dealing with films deposited by unbalanced reactive magnetron sputtering, three different structures of CN compounds have been identified depending on the deposition temperature and nitrogen concentration. These structures include an amorphous phase which forms at low temperatures (~100°C), a graphite-like phase which forms at low nitrogen content and high temperatures (> 200°C) and a fullerene-like structure for high temperatures (> 200°C) and nitrogen concentrations above 10 at%. An extension of that phase diagram for ta-C films deposited by FCVA has been done by Kleinsorge et al. (2000). In this case, doping the ta-C films can only be performed at deposition temperatures below 100 °C and nitrogen concentrations below 0.4 at%. Increasing the temperature or the nitrogen content up to 9 at% induces a sp2 clustering although the sp3 content remains inalterable. However, increasing the deposition temperature above 200 °C leads to a graphitic film, whereas nitrogen contents above 9 at% leads to a sp2 bonded matrix transformation, where fullerene-like structures can be observed. As the formation of C3N4 requires a 100% sp3 bonding coordination for carbon and sp2 for nitrogen, there was a considerable interest in obtaining the highest possible incorporation of nitrogen as well as the highest fraction of sp3 bonded carbon. The main techniques used to characterize the chemical structure of the deposited films have been XPS and XAS/EELS combined with other complementary techniques (e.g., IR, Raman, NMR, etc.). In any case, there have been considerable problems in establishing an unambiguous local bonding configuration of the deposited films. The XPS data was rather controversial (Muhl and Mendez, 1999; Hellgrenet al., 1999; Cheng et al., 2000; Quiros et al., 2000) with respect to the assignments of the spectral features to particular bonding structures. Nevertheless, some general agreement was found on the assignment of the two main components observed in the Nls XPS spectrum. They were associated with nitrogen bonded to sp2 and to sp3 coordinated carbons atoms. A third component commonly observed in films with high nitrogen content, is attributed to nitrile groups. In general, the Cls XPS spectrum appears rather broad and featureless, so that its analysis in terms of sp2/sp3 carbon is ambiguous and controversial. To avoid this ambiguity, many studies use XAS or EELS to measure the 7i-states in the Cls absorption edge as an estimation of sp2 and sp1 hybridised carbon (Hellgren et al., 1999; Quiros et al., 2000). Raman spectroscopy has also been widely used, as in the case of DLC films. The spectra present the common G
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and D bands (cf., section 5.2.3), which are usually analysed in terms of intensities and widths as a measure of the order and the size of the graphitic domains in the amorphous film (Hellgren et al., 1999; Cheng et al., 2000; Muhl and Mendez, 1999). From the corresponding plasmon loss energy (cf., section 5.2.4), the density of CNX films deposited by IBAD has been estimated to be around 2.2-2.5 g cm"3 Other properties of amorphous CNX films have also been evaluated and extensively reviewed (Wang, 1999). The mechanical properties are comparable with those of the amorphous DLC. In fact, deposition of CNX films at temperatures above 200 °C leads to graphite-like films with very poor mechanical properties. However, carbon nitride films grown at low temperatures, with a nitrogen content around 15-20 at %, have shown interesting properties and potential applications. Apparently these films combine high hardness with an extreme elasticity and very smooth surface. Some authors suggest that it is due to a fullerene-like microstructure, which is believed to be promoted by the incorporation of nitrogen. Elastic recoveries of up to 90% and hardness between 20-28 GPa. have been reported by nano-indentation for samples made by different techniques. The smooth surfaces of these films, with friction coefficients as low as 0.1-0.2 under different experimental conditions (Donnet et al., 1999), make these films also appropriate for tribological applications in hard discs of high data storage density. The stress of the films has been measured to be compressive in all the cases. Films deposited by unbalanced magnetron sputtering show a relatively low stress around 1-1.5 GPa independent on the nitrogen content. However those films grown at higher temperatures (350-550°C) show a rapid increase of the initially very low stress (near cero) as the nitrogen content is above 5 at % (Hellgren et al., 1999). In the case of ta-C films doped with nitrogen, there is a continuous stress relaxation as nitrogen incorporates or as the nitrogen ion energy is increased, a feature that has been attributed to the increase of both the sp2 fraction and the defect density in the film (Cheng et al., 2000). The search for new hard materials with pure covalent bonds, where the formation of sp3 hybridisation is of major importance, has focused a lot of effort in the last years. In this context, the system B-C-N seems to involve promising candidates. A schema of the B-C-N ternary phase diagram is shown in Figure 5.17. Main binary compounds have been included in the corresponding sides while the ternary stoichiometric BC2N and BCN have also been included. Ternary BCN cubic structures are expected to combine the properties of c-BN and diamond, e.g.,
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hardness between 40 and 90 GPa and high thermal stability. The system has been the subject of both theoretical (Tateyama et al., 1997; Mattesini and Matar, 2001) and experimental work (Knittle et al., 1995; Solozhenko et al., 2001) and since the super-hard cubic BC2N phase has been synthesised under high T and high P conditions (Solozhenko et al., 2001) the motivation for its synthesis under metastable conditions has increased considerably. Experimentally, BCN materials with different compositions have been produced by a large variety of low pressure methods (CVD and PVD), although its characterization is still scarce. In general only hexagonal soft phases have been obtained, although many efforts have been made to induce the highly dense cubic phases.
Figure 5.17. B-C-N diagram including some relevant compounds.
IAD and IBAD methods have been used to deposit BCN compounds by evaporation, ablation or sputtering of different compounds (B4C, BN + graphite) assisting the film growth with several mixtures of gases (e.g., N2+, N2+ + Ar+, N2+ + Ar+ + CHt-1", etc.) and under different experimental conditions (Yap et al., 2001, Gago et al., 2002). In general only soft, hexagonal ternary (BCxNy) or mixtures of segregated phases (CNX + BNX + BXC) are obtained. Interestingly, Yap et al. (2001) report the deposition of BC2N on Ni at 800°C by assisted PLD of a graphite + BN target. These authors claim that the Ni acts as a sink for carbon at high temperature and therefore prevents the formation of graphite, so that a ternary BCN solution
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with composition BC2N is obtained. However, the properties of the films were not published. Ternary compounds obtained by evaporation of B4C assisted with a gas mixture of N2+ + Ar+ + CU/ have been characterised by Gago et al. (2002). A hardness of around 35 GPa and friction coefficients of -0.05 were measured for thin films of that material deposited under optimal conditions.
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Solozhenko, V.L., Dub, S.N., Novikov, N.V., DiamondRel. Mat., 10 (2001) 2228. Spencer, E.G., Schmidt, P.H., Joy, D.C., Sansalone, F.J., Appl. Phys. Lett., 29 (1976)118. Uhlmann, S., Frauenheim, Th., Lifshitz, Y., Phys. Rev. Lett. 81 (1998) 641. Uhlmann, S., Frauenheim, Th., Stephan, U., Phys. Rev. B 53 (1996) 4117. Wang, E.-G., Adv. Mater. 11 (1999) 1129. Voevodin, A.A., Donley, M.S., Surf. Coat. Technol. 82 (1996) 199. Watanabe, S., Miyake S., Murakawa, M., Surf. Coat. Technol. 49 (1991) 406. Weiler, M., Sattel, S., Giessen, T., Jung, K., Ehrhardt, H., Phys Rev. B 53 (1996) 1594. Weissmantel, C , J. Vac. Sci. Technol. 18 (1981) 179. Weissmantel, S., Reisse, G., Thin Solid Films 355-356 (1999) 256. Windischmann, H., J. Appl. Phys., 62 (1987) 1800. Xu, S., Tay, B.T., Tan, S., Zhong, L., Tu, Y.Q., Silva, S.R.P., Milne, W., J. Appl. Phys., 79 (1996) 7234. Yap, Y.K., Wada, Y., Yamaoka, M., Yoshimura, M., Mori, M., Sasaki, T., Diamond Rel. Mat. 10(2001)1137. Yoshikawa, M., Ishida, H., Ishitani, A., Kouzumi, S., Inuzuka, T., Appl. Phys. Lett. 58(1991)1387. Zellama, K., Current Opinion in Solid State and Materials Science 4 (1999) 34. Ziegler, H., Schreibe, H.J., Schultrich, N.B., Surf. Coat. Technol. 97 (1997) 385.
ACRONYMS LIST
a-C
Amorphous carbon
a-C:H
Hydrogenated amorphous carbon
AES
Auger electron spectroscopy
AFM
Atomic force microscopy
AZO
Aluminum zinc oxide
CIB
Cluster ion beam
CVD
Chemical vapour deposition
DC
Direct current
DIBS
Dual ion beam sputtering deposition
DLC
Diamond like carbon
DSFR
Dynamic scaling function of roughness
EBD
Electron beam deposition
ECR
Electron cyclotron resonance
ECWR
Electron cyclotron wave resonance
EELS
Electron energy loss spectroscopy
ERD
Elastic recoil dispersion
EXAFS
Extended X-ray absorption fine structure
FDLC
Fluor-DLC
FED
Field emission displays
FCVA
Filtered cathodic vacuum arc
FCVAD
Filtered cathodic vacuum arc deposition
FET
Field emission transistor
FTIR
Fourier transform infrared
FVAD
Filtered vacuum arc deposition 275
276
Low ENERGY ION ASSISTED FILM GROWTH
GCIB
Gas cluster ion beam
GMR
Giant magnetoresistance
HAP
Hydroxyapatite
HOPG
Highly oriented pyrolitic graphite
I/A
Ion to atom ratio
IAAD
Ion assisted arc deposition
IAD
Ion assisted deposition
IAPLD
Ion assisted pulse laser deposition
IAPVD
Ion assisted physical vapor deposition
IB AD
Ion beam assisted deposition
IBD
Ion beam deposition
IBICVD
Ion beam induced chemical vapour deposition
ICB
Ionised cluster beam
ICP
Inductively coupled plasma
IMS
Ionised magnetron sputtering
IMPVD
Ionised metal physical vapour deposition
IP
Ion plating
ITO
Indium tin oxide
LTI-
Low temperature isotropic (LTI) pyrolitic carbon
MBE
Molecular beam expitay
MC
Montecarlo
MD
Molecular dynamics
MS
Magnetron sputtering
MSIBD
Mass selected ion beam deposition
NMR
Nuclear magnetic resonance
NRA
Nuclear reaction analysis
PACVD
Plasma assisted CVD
ACRONYMS LIST
PBS
Plasma beam source
PUD
Plasma immersion ion deposition
PHI
Plasma immersion ion implantation
PVD
Physical vapour deposition
RBS
Rutherford back-scattering spectroscopy
RF
Radio frequency
RHEED
Reflected high energy electron diffraction
RICBD
Reactive ion cluster beam deposition
RMS
Root mean square
STM
Scanning tunnelling microscopy
ta-C
Tetrahedral amorphous carbon
ta-C: H
Hydrogenated tetrahedral amorphous carbon
TCO
Transparent conductive oxide
TEM
Transmission electron microscopy
TRIM
Transport of ions in matter
UHV
Ultrahigh vacuum
ULSI
Ultra large scale integration
UV-Raman
Ultraviolet-Raman
Vis-Raman
Visible-Raman
XANES
X-ray absorption near edge structure
XAS
X-ray absorption spectroscopy
XPS
X-ray photoelectron spectroscopy
XRD
X-ray diffraction
XRR
X-ray reflectometry
YBCO
YBa2Cu307.5
YSZ
Yttria stabilised zirconia
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SUBJECT INDEX
Deposition methods Deposition parameters Stress Properties Tribological Optical Electronic C-B-N C/Si thin films Child-Langmuir state Coatings Corrosion resistant
Accelerated neutral molecules, 214 Effects on thin films 92 Adatoms 100 Mobility 114 Mobility and densification Adhesion 108,164 Interface mixing 166 Polymer-metal interface 126 Amorphisation 160 Atom peening 95 Au particles, nucleation Ballistic interactions Biaxial oriented films Cr/CoCrPt films Binary elastic collisions Binding energy Surface binding energy Lattice binding energy Biomaterials Modification Corrosion protection Biocompatibility Bragg's law Bradley's model for texture Broad beam ion sources Carbon nitride CNx Cascade formation c-BN Characterization XRD FTIR EELS/XAS Cross section
90 143,149 274 12 4 5 5 186 186 186 187 137 148 80
Hard Metal Nitride Oxide Optical Solid lubricants Tribological Wear resistant Columnar growth Compound formation by IAD diagram Columnar growth DLC c-BN nucleation C-B-N Cr-N phase diagram Coordination number Corrosion Resistant coatings Biomaterials Magnesium alloys Zinc, zinc alloys
265 34 245 246 247 247 248 250
279
251 252 255 ,260 256 256 257 257 264 111 77 179, 183, 184 ,186 175 181 181 182 190 177 174 175,185 93 ,113 130 113 218 254 268 134 4 179 186 183 184
280
Low ENERGY ION ASSISTED FILM GROWTH
Coulomb potential CrN and Cr2N synthesis Crystallisation Effect of temperature Effect of ion assistance Ti0 2 thin films ZnO thin films AI2O3 thin films
8 216 126 127 128 128 128
Damage Average depth of damage 26 30 Displacement damage function Efficiency 28 27 Ion bombardment 163 Davis's model 119 Defects in IAD thin films 122 CrNx films 120 Surface and bulk 122 Thin film microstructure 119 TiN thin films 122 XRD monitoring of defects 114 Densification 116 Density and crystallinity Subplantation model 118 ,263 218 Diamond like carbon 218 a-C 218 a-C:H 218 ta-C 218 ta-C:H 221 Atomic structure Characterization 220 228 Cross section 225 Density 228 Deposition methods 231 Deposition parameters Ion energy 231 235 Substrate temperature 238 Other parameters
221 Electron scattering Energy loss 223 Plasmon energy 225 Properties 241 Electronic 243 242 Mechanical Optical 243 242 Tribological 222 Sp3 bonding fraction 222 Sp2 bonding fraction Stress 239,260 60 Direct ion assisted deposition 3,28 Displacement energy 147 Dobrev's model 55 Dual Ion Beam Deposition 61 Electron Cyclotron Resonance 51 Electron evaporation source 19 Electronic energy loss parameter 84 End-Hall ion sources Energy 25 Depth distribution function 5 Displacement 41 Normalized Reduced 11 49 Ranges 5 Thermal Transferred in a single collision 14 Epitaxial thin films 70 ,108 108 GaN Evaporation 49 Filament-less ion sources Filtered vacuum arc deposition Frenkel pair Fretting damage Fretting wear Fullerenes
86 66 5 185 185 72
SUBJECT INDEX
GaN epitaxial films Gas adsorption Glow discharge plasma Graded composition Grain size Growth mechanisms
108 52 64 214 99 92
Hard coatings Heat of sublimation Hollow cathode ion source Hydroxyapatite
175 5 86 186
123 Inert gas incorporation 124 Ar in ZrC>2 Compressive stress 124 ,161 Interface mixing 108 Mixing by high energy ions 109 Mixing in IAD thin films 110 Mixing efficiency function 111 Ion Assisted methods 47 72 Beam deposition Beam induced chemical vapour deposition 58 14 Energy loss rate Implantation 76 Ion to atom arrival ratio 42 44 Momentum transfer Plating 60 Sources 80 Velocity 8 Ionised cluster beam 69 Ionised magnetron sputtering 64 Ionised metal PVD 65 Kaufmann ion sources
81
Laser ablation
52
Lennard-Jonnes potential Lindhard-Scharff model Lubricant coatings
281 4 19 177
Magnesium alloys corrosion 183 204 Magnetic films 205 Giant magnetoresistive 207 Hard bias 206 Magnetic heads 205 Magnetoresistive 204 Metallic Magnetron sputtering 64 Mass selected Ion Beam deposition 72 Metallic coating 181 ,204 187 Metallisation of polymers 133 Metastable phases 5 Miller indexes 258 Models of growth of DLC 259 Preferential sputtering 260 Stress models 261 Thermal spike 262 Subplantation Molecular dynamics simulations 105 110 Mixing efficiency 111 Void removal 148 Texturing 197 Narrow band filters 182 Nitride coatings Norgett-Robinson-Torrens model 28 92 Nucleation 96 Ni films 251 c-BN films Optical coatings Dielectric oxide Fluoride Narrow band filters
190 191 196 197
282
Low ENERGY ION ASSISTED FILM GROWTH
Rugate filters Transparent conducting Oxide coatings
199 201 182 ,191
Penning ionisation 62 Phase diagram 134 Cr-N 218 DLC c-BN nucleation 254 CBN 268 Planarisation of surfaces 103 Plasma Discharge 61 Immersion ion deposition, PUD 78 Immersion ion implantation, PHI 76 Sheath 61,77 94 Percolation of particles 141 Polar plots 187 Polymers Potential 4 Lennard-Jonnes Thomas Fermi 8 10 Power law approximation 137 Preferential orientation Primary knock on atoms (PKA) 27 PVD 92 Range Total Projected Reactive ion cluster beam Reflected high energy electron diffraction RF filament-less ion source Rocking curves Roughness Interfaces Scaling theory Sputtering
20 22 72 72 87 140 98 99 97 101
Surfaces Rugate filters Rutherford Backscattering Spectroscopy Scattering Elastic Nuclear Electronic Screening Function Length Smoothing of surfaces Spike formation Sputtered atoms Angular distribution Energy distribution Sputtering Sputtering yield Step coverage Sticking probability Straggling Range straggling Energy straggling Stopping Cross section Electronic Nuclear power Sub-surface implantation Subplantation model Stoney's formula Strain Stress c-BN films Compressive Cr films DLC films
99 199 12,17
12 12 12 9 9 100 30 39 40 36 37 65, 103 72 25 27 15 19 16 15 74 263 155 153 255, 260 160 159 239, 260
283
SUBJECT INDEX
IAD thin films Intrinsic PVD thin films Tensile Thermal Surface defects Ta 2 0 5 thin films Temperature Local Equivalent Texture Coefficient IAD thin films PVD thin films Oxide thin films Nitride thin films Applications Thermal energy Thermal spike
105, 133 130 67 49, 174
153, 158 156 156 160 155 94
TiN thin films TiC>2 thin films Toroidal magnetic field Tribological coatings
102
Wear resistant coatings Windischmann's model
175 159, 162
XANES and amorphisation X-ray reflectometry XRD
130 106, 225 137,,248
35 49 138 142 142 144 145 152 6 32
66
Vacuum arc
YSZ biaxially aligned thin films YSZ buffer layers
150 152
Zinc alloys corrosion ZrN and Zr3N4
184 135
