THE STUDY OF DEFECTS IN LOW MISFIT GE-SI

STRAINED LAYER HETEROSTRUCTURES

Mark Benedetto Stirpe

A thesis submined in conformity with the requirements for the degree of Master of Applied Science

Graduate Department of Metallurgy and Materials Science University of Toronto

0 Copyright by Mark Benedetto Stirpe 1998

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The Study of Defmts in Low Misfit Ge-Si Sfrained Layer Heterostructures

Master of Applied Science Mark Benedetto Stirpe

Graduate Department of Metaiiurgy and Materials Science University of Toronto

In order to study the defect structures in strained layer superlattices, a series of

Ge,SiI.,/Si (100) (0.09 c x < 0.13) heterostructures were grown by u l a g h vacuum

chernical vapour deposition (UHV-CVD) and molecular beam epitaxy (MBE) methods.

Activation energies were determined to be consistent with the widely accepted values (Q,,

= 2.5 + 0.5 eV and Q, = 4.2 + 0.5 eV for nucleation of misfit dislocations and overd1

suain relaxation, respectively). The low-temperature MBE growths yielded Q,, = 0.5 k

0.05 eV and Q, = 2.0 f 0.5 eV for nucleation and overall strain relaxation, respectively.

Samples implanted with Si ions displayed significant decreases in nucleation rate and

misfit dislocation densities. Based on the quantitative study conducted on GexSiiJSi

heterostructures, it has been demonstrated that the onset of strain reIaxation via misfit

dislocations can be controlled by point defcct injection via ion implantation.

I would like to thank my supervisor, Professor D. D. Perovic for his advice and

encouragement throughout the course of this thesis. 1 wodd also like to thank Dr. H.

Lafontaine, Dr. J.-M. Baribeau, and Dr. D. C. Houghton at the National Research Council

of Canada (NRCC) for participating in useful discussions and providing excellent

samples for study. Speciai thanks also go out to Dr. R Goldberg, Dr. P. Simpson, and

Dr. 1. Mitchell at the University of Western Ontario for the ion implantations and positron

annihilation spectroscopy.

1 wodd also like to extend my appreciation to Mr. Sd Boccia, Mr. Fred Neub, Ms.

Azita Ariapour, and Mr. Bahi Bahierathan at the University of Toronto for their help with

sample preparation and TEM instruction.

Finaily, 1 am grateful to the N a W Sciences and Engineering Research Councii of

Canada and the University of Toronto Fellowship Award for financial support and the

National Research Council of Canada for travel suppoa through the Visiting Research

Graduate Program.

TABLE OF CONTENTS

Page #

Abstract

Acknowledgements

Table of Contents

List of Figures

List of Tables

Acronyms

Chap ter 1- Introduction

1 - 1 Ge-S i Heterostructures

1.2 Molecular Beam Epitaxy

1 -3 Chernical Vapour Deposition

1.4 Ion Implantation

1.5 Positron b h i l a t i o n Spectroscopy

1 -6 Microscopy Techniques

1.6.1 Nomarski Interference Microscope (NIM)

1 A.2 Field Ernission Scanning Electron Microscope (FE-SEM)

1 -6.3 Transmission Electron Microscope (TEM)

Chap ter 2 - Experimental

2.1 Heterostmcture Growth Conditions

2.2 Ion Implantation Conditions

2.3 Rapid Thermal Annealhg (RTA)

2.4 Nomarski Interference Microscopy (NIM)

2.5 Field Emission Scanning Electron Microscopy (FE-SEM)

2.6 Transmission EIectron Microscopy (TEM)

iii

2.7 Positron Annihilation Conditions

Chapter 3 - Results and Discussion

3.1 Strained Layer Geometry

3 -2 Buik Quantitative Measurements

3.3 Sources of Nucleation

3.4 Ion implantation

3.5 Low-Temperature MBE Heterostnictures

Chapter 4 - Conclusions and Recommendations

Appendix

Effective Stress and Kinetic Mode1

Energy Balance Approach

Bulk Quantitative Raw Data

TRlM Simulation

References

LIST OF FIGURES

Figure 1.1.1 - Heteroj unction Bipolar Transistor

Figure 1.1 -2 - Diagram of alternative modes of epitaxial growth

Figure 1.1.3 - Kinetically Iimited 'critical' thickness curves

Figure 1.1.4 - Schematic diagram of prismatic loops at Ge-rich platelets, half-loops and expansion of lwps

Figure 1 -2.1 - Schematic diagram of MBE process

Figure 1.2.2 - Plot of epitaxiai film morphology as a function of growth temperature and Ge composition for t 000 A Ge&,

Figure 1 -3.1 - Schematic diagram of UHV-CVD process

Figure 1 -4.1 - Photo of typical Ion bearn implantation device

Figure 1.5.1 - Schernatic diagram of positron's reaction in solids

Figure 1 S.2 - (a) Doppler broadening (b) positron lifetimes, and (c) S-parameter

Figure 1.5.3 - Schematic diagram of back reflection geometry in positron beam production

Figure 1.6.1 - Nomarski interference microscope

Figure 1.6.2 - FE-SEM gun configuration

Figure 1.6.3 - Relationship between total emitted e- coefficient and beam energy

Figure 2.1.1 - Czochralski growth method

Figure 2.4.1 - Typicai Nomarsici micrograph

Figure 2.6.1 - Schematic process of TEM cross-section preparation

Figure 2.6.2 - (a) Schematic of difhction pattern of diamond cubic lattice with beam direction z = [O0 11 (b) Actuai TEM image showing two-barn condition g = [400] (c) Schematic of cross-section conditions with z = [O 1 11

Figure 3.1.1 - Geornetncal distribution in strained layer superlattice

Figure 36.1 - Nomarski images for temperatures ranging h m 700 - 1 0OO0C

Figure 3.2.2 - Nucleation rate vs. F' of UHV-CVD and MBE matenal

Figure 3 -2.3 - Misfit dislocation density vs. K" for UHV-CVD and MBE Matenal

Figure 3.2.4 - Nucleation rate vs. K" for UHV-CM as-gmwn and implanted material

Figure 3 .ZS - Nomarski micrograph showing as-grown and ïmplanted sides after RTA at 850°C for 30 s

Figure 3.2.6 - Misfit dislocation density vs. EC1 for CVD-61 as-grown and irnplanted matenal

Figure 3.2.7 - Nucleation rate vs. Kf for UHV-CVD and low-T MBE Material

Figure 3.2.8 - Misfit dislocation density vs. K" for UHV-CVD and low-T MBE matenal

Figure 3.2.9 - Positron annihilation spectroscopy for Iow-T MBE material And CVD-9

Figure 3.3.1 - TEM DF image Ge-rich pïatelets, g = [220]

Figure 3.3.2 - TEM BF image of C M - 9 after RTA for 5 s at 1000°C

- l q - l Figure J .J -3 - TEM DF image of CVD-9 d e r RTA for 5s at 1 OOO°C

Figure 3.4- 1 - TEM BF image of implantation profile in cross-section with z = [O1 11 for RTA CVD-9 at 850°C for 30 s

Figure 3.4.2 - TEM plan-view image of implanted CVD-9 RTA at 850°C for 30 s

Figure 3.5.1 - Morphological instability phase diagram for Ge&30.74Si

Figure 3-52 - FE-SEM SE-detector image showing presence of voids

Figure 3.5.3 - TEM cross-section image revealing cusps wiîh ( 1 1 1 1- oriented facets

Figure 3 -5.4 - Doppler-broadened annihilation line-shape parameter S as a function of positron energy

LIST OF TABLES

Table 2.1.1 (a) Growth parameters of UHV-CVD G&Si 1JSi materiai

Table 2.1.1 (b) Growth parameten of MBE Ge&.,JSi matenai

Table 2.3.1 (b) Rapid Thermal Anneai TemperatundTime parameters

HBT - heterojunction bipolar transistor

IC - integrated circuit

LED - light - emitting diode

MBE - molecular b a r n epitaxy

UHV-CVD - ultrahigh vacuum chemical vapour deposition

TEM - transmission electron microscope

No - inifial misfii source density present in as-grown material

Q. - activation energy for nucleation of misfit dislocations

Q, - activation energy for overall h n relaxation

RTA - rapid thermal anneai(s)

LPCVD - low pressure chemical vapour deposition

sccm - standard cubic centimeten per minute

AES - Auger electron spectroscopy

XRD - X-ray diffraction

PL - photoluminescence

SIMS - secondary ion rnass spectroscopy

LIBI - linear ion beam implantation

y - ray - gamma rays produced in positron annihilation spectroscopy

NIM - Nomarski interference microscopy

FE-SEM - field emission scanning electron microscope

BS - backscattered electrons

SE - secondary electrons

DP - diffraction pattern

z - beam direction

BF - bright field

DF - dark field

SLS - strained layer superlattice

RTCVD - rapid thermal chemical vapour deposition

Er - energy of vacancy formation

viii

E, - energy of vacancy migration

Low-T - Iow-temperature

TRIM - transport of ions in matter

BCA - binary collision approximation

MD - molecular dynarnics

CHAPTER 1 - INTRODUCTION

1.1 Ge-Si Heterostnictures

Within the past decade, techniques for growing crystalline materiais have advanced to

the point whexe germanium can be selectively introduced into silicon. Germanium is

grouped with silicon in the periodic table and possesses the same diamond cubic crystal

structure and the same bonding orbitais (Le. four tetragonal sp3 hybnds). When

germanium is doped to the base region of a silicon bipolar transistor, the result is a Si-Ge

heterojunction bipolar transistor (HBTs)(Figure 1.1 -1). The device properties can be

tailoreci tightly to the needs of the desired application, because the profile and

concentration of the germanium can be wntrolled with great accuracy. Hence the IC

designer's new-found ability to devise hi&-speed, lowast chips using heterostruchired

layers (Cressler, 1 995).

Figure 1.1-1:

Al-CU

n' polysilicon Silicon-Germanium

P+

Schematic d i a m of a heterojunction bipolar transistor (HBT) in an advanced SiGe IC

The HBT is present in a self-aligned SiGe IC. Each HBT is isolated by I .O pm-wide

trenches etched 4.0 - 5.0 pn into the silicon subsîrate and overgrown with a composite of

polysilicon and oxide. Heavily doped d and P* polysilicon layers form the contacts to

the emitter and base legions, respectively. The AI-Cu is used for electrical contacts.

Due to the tremendous advancement in Si-based technology, considerable attention

has been given to the development of strained layer heterostruchues consisting of the

group-IV alloys Si-Ge. The retention of structurai perfection during epitaxial growdi and

thermal processing is crucial for the development of GqSil,/Si seained layer devices

such as HBTs, resonant tunuehg diodes, and Light-emitting diodes (LED). Such

heterostructures are, in generd, metastable and can relax through the injection of misfit

dislocations at the Ge,&,/Si interfaces upon elevated temperature exposure (Houghton,

1991). Misfit dislocation propagation must be suppressed to avoid the shultaneous

formation of threading dislocations, which may penetnite heterojunctions and enhance

current leakage. Successful use of semiconductor heterostructures in electronic and

optoelectronic devices depends strongly on the reduction of dislocation density below 10'

cm-2 (Rajan er al., 1991). The 4.2% lattice mismatch that exists between Si and Ge is

enough to cause this disruption in the crystalline order at the GeSVSi interface.

The concept of strained layer epitaxy was considered fim by Frank and van der

Mewe (1 949) and then M e r developed by Matthews and CO-workers (1972). Epitaxy

can be defined as an oriented overgrowth in which each successive layer tends to assume

the lattice orientation of the layer beneath it (Perovic, 1988). in the case of G ~ s i ~ . ~ / S i

heterosuucnires, the Larger lattice parameter of the epilayer, GeSi, must elastically

cornpress biaxially in the interfacial plane to match the lattice parameter of the substrate

(Le. Si). The misfit strain is then accommodated by a tetragonal distortion of the growllig

epilayer in the growth direction normal to the interface. Figure 1.1.2 shows a diagram of

the pmcess. The formation of misfit dislocations at the interface relieve this strah. The

misfit dislocations, typically 60' al2 4 10> in character, are parallel to the interfaces and

the thrmding dislocations gliding on the { 1 1 1 } -variants are the geometries expected in

partially relaxed heterostructures (Houghton, 199 1).

EpitaxiaI layer

Substrate crystal

(a) unsmined

(b) strained

Figure 1.1.2: Epitaxial growth that gives (a) an unstrained layer with misfit dislocations at the interface and (b) a straïned layer which deforms to match the lanice spacing

There exists a 'critical thickness' for strained layes where coherent dislocation-fke

growth takes place. The primary components for such growth are the control of the Ge

composition and thicknesses of layers. Figure 1.13 shows the kineticaily limited cntical

thickness of previous studies in Gedi i,/Si research at various temperatures. This criticai

thickness detemiines the metastability limiîs for strained layers. Layers grown above the

cntical thickness will nucleate interfacial misfit dislocations. Layea grown below the

critical thic kness will not. However, metastable strained layers can be produced which

are thicker than their equilibrïurn critical thickness. The metastability l imits are

determined kinetically by the availability of thermally activateci nucleation sites for misfit

dislocations and their rate of extension by glide. and by growth parameters such as

temperature and deposition rate (Houghtoh 1995). Growth techniques such as low

temperature rnolecular beam epitavy (MBE) and reduced temperature chemical vapour

deposition (CVD) are utilized to develop these structures. Under proper groowrh

conditions, strained layer epitavy allows one to repeatedly grow Ge,Sii, strained layers

separated by Si layen so that the whole layer sequence maintains the same lanice

parameter as the Si substrate.

10 20 30 40 50 60 Ge concentration (% 1

Figure 1 -3.1 : Kineticall y limited "cri tical" thickness for Ge,Si &i (1 00) growth

In studying strained layer heterostructures, it is also of importance to consider the

structurai limitations of growth techniques; in particular, low-temperature Si-MBE

growth. Growth at lower temperatures is desired for increasing the thicknesses of layers

while still maintahhg coherency among the layers. However, void formation cm be

associated wiîh low temperature MBE (<400°C). A study of void formation has been

conducted and will be discussed.

There have been rnany different treatments of the critical thickness model in the past,

but the treatment of Matthews and CO-workers has proved to be the most popular

(Perovic, 1995). The theory of Matthews et d. (1970) takes into consideration pre-

existing sources of nucleation for misfit dislocations such as substrate threading

dislocations. By takllig this into account, one can relate the equilibrium critical thickness

to the first misfit dislocation that is inîmduced at a misfihg interface (Perovic et al.,

1995).

In order to study the s& relaxation effects in metastable strained layer

heterostnictures, non-equi librium treatments of the pro blem would have to be

investigated. People and Bean (1985) aîtempted to describe the critical thiclmcss in

GeSi/Si using an energy balance model that considered spontaneous homogeneous

interfacial nucleation of misfit dislocations in the absence of pre-existing dislocations.

There are kinetic (i.e. Peieris) barriers, which iimit the extent of s& relaxation in

metastable structures for given temperature and tirne. Therefore, although exceeding the

critical thickness will result in misfit dislocations, it is not sunicient in determinhg the

onset of strain relaxation. Therefore, it is not possible to compare metastable structures

with equilibriurn theory (Perovic et al., 1995).

To smdy the effects of seain relaxation on sttained layer heterostructures, kinetic

barriers would have to be eliminated via post-growth annealing, and methods of detecting

thefint stages of misfit dislocation would have to be implemented. An excellent method

for determining the onset of strain relaxation is through the use of chernical

etching/lrlomarski interference microscopy, which will be discussed in later sections.

This technique is able to detect c 104 dislocations/cm2 unlike X-ray difiction, TEM, ion

channeling, Raman spectmscopy, photoliIminescence spectroscopy, and laser light

scattering. Houghton et ai. (1990) and Zhao (1993) were able to use Nomarski

microscopy to obtain results in excellent agreement with equilibriurn predictions.

The sources of nucleation of misfit dislocations were extensively exarnined by the

group of Perovic et ai. (1990). It was show using transmission electron microscopy

(TEM) that interfacial nucleation of 60' misfit dislocations was associateci with sub-

nanometer sized Ge-rich platelets that evolve fiom strain-induced growth (Perovic and

Houghton 1992, 1993). Rowell et al. (1993) found that the pIatelet formation has been

attributed to the onset of 2-D islanding under step growth conditions. The Ge

preferentially accumulates at the periphery of atornic terrace kink sites and is then

covered by the next atomic Iayer during the growth procedure (Perovic et ai. (1995).

This study M e r considered the nucleation processes in low misfit heterostnicnires of

Ge,SilJSi, where xc 0.13.

Low misfit heterostnicnires were chosen for this study to maintain consistency with

the nucleating mechanisms taking place Ï n the "Stage I" regime (dislocation densities

1 O' cd2) . Perovic and Houghton (1 995) defined this regime to be nucleation-limited,

where misfit dislocation nucleation was observed to increase Linearly with tirne, beyond

some background density, No (see Appendix). The misfit dislocation velocity remains

constant through this regime resulting in a hear increase in the oved l strain relaxation

rate with time. Activation energies were observed to be Q, = 2.5 + 0.5 eV and Qr = 2.25

f 0.05 eV for nucleation and glide of misfit dislocations, respectively (Houghton, 1991).

The activation energy for strain relaxation in the Stage 1 regime was fomd to be Q, =

4.75 f 0.5 eV. This value is consistent with combining the nucleation and glide energies

(Le. Qr = Qn + Qv) in series.

Interfacial perturbations

( 100) Substrate

Figure 1.1.4: Schematic diagram illustrating the nucleation of prismatic dislocation loops (vacancy-type) at the Ge-ptatelets localized at the strained layer interface. The prismatic loops generate the haif-loops on { 1 1 1 }variants which then expand to forrn misfit dislocations with an extemal driving force (RTA).

As in the work done by Perovic and Houghton (1992, 1993), the nucleation of misfit

dislocations of the strained layer heterostructures of this study will be considered as a

two-step process. The fmt step is the nucleation of the initial pnsmatic dislocation loop

(Le. with the loss of coherency at the Ge interstitial perturbation) and the second step is

the expansion of the loop to fonn a stable misfit dislocation. Figure 1.1.4 shows a

schematic diagram of the pmcess. The prismatic loops generate 60' shear loops, which

form half-loops in the epilayer andor multiple loops in the bufEer layer. With the aid of

some extemal driving force (Le. rapid themial annealhg (RTA)), the half-loops expand

leaving behind 60' misfit dislocations in two orthogonal 4 10> directions.

1.2 Molecular Beam Epitaxv

Ge,Sii., heterostnicnires can be grown at low temperatures (350-700°C) using the

MBE process. The advantages of p w i n g heterostructures at low temperatures is to

reduce the vibrational energies of thé growing film to decrease the disorder and

interdiffusion of Ge in Si (Hull et al., 1985). Figure 1.2.1 shows a schematic illustration

of the process. MBE is a technique for growing overlayers on crystals by the

condensation of low-vapour-pressure molecules or atoms from a molecular beam in a

high vacuum. The low vapour pressure of the materiai is required to provide nearly

instantaneous control of the deposition fluxes. Decomposition of the molecules at the

surface is avoided by using simple molecules or atoms.

O Heated rotating substrate crystal

Doping Gun

Figure 1.2.1 : Schematic diagram of MBE process

8

Si-based MBE consists of the coevaporation of sficon and dopants and acceleration

onto a slightly heated substrate. Silicon and germanium condense once they reach the

surface due to their negligible room temperature vapour pressures. This method of

elemental evaporation p e d t s the growth of crystalline layers at temperatures where

solid-state diffusion is negligible. Suice chernical decomposition is not required,

deposition species need acquire only enough energy to migrate dong the substrate

surface to a cry stalline bonding site (Bean, 1 98 1). This requires only minimal substrate

heating.

A Czochralski-grown silicon d e r of suitable orientation (i.e. <100>) is placed

polished side down in an ultra-high vacuum chamber. An in-situ cleaning technique is

used to produce an atornically ciean surface pnor to growth. Several methods can be

used such as ~ r ' sputtering/annealing and thermal desorption (1200°C) (Perovic, 1988).

Afier the substrate wafer has been cleaned, it is heated to its desired growth temperature.

The heated substrate ensures that single crystal growth will take place because the

thermal energy produced increases the sdace migration energy of atoms so they reach

proper bonding sites at the fke sdace. Deposition rates and concentrations are

controlled by rehctory metal shutters which are moved into and out of the evaporatiun

path. This process is cornputer controlled.

If proper conditions are met, smooth 2-D growth will take place. Lower growth

temperatures and lower Ge concentrations will ensure 2-D growth. However, if the

temperature and Ge concentration are raised accordingiy, 3-D growth will take place.

The kinetic pathways of the 2-D-to-343 transition are detailed by the group of Jesson et

al. (1996). Figure 1.2.2 displays a plot of epitaxial film morphology as a bction of

growth temperature and Ge concentration for a 1000 A thick GexSil, füm. Several in-

situ techniques are used to monitor the epitaxial growth parameters. A quartz crystd

deposition meter has a fmluency which changes during excitation as the crystal thickens.

Growth rates as low as - 1 A/sec can be realized with an accuracy of approximately 10%.

As well, there is a quadraple mass spectrometer which can be used to obtain molecuiar

beam flux and residual gas analyses. The surface crystallinity of the growïng crystal can

also be monitored. A surface electron difiction technique such as high energy electron

diffraction or low energy electron di f ic t ion can be utilized (Perovic, 1 98 8).

Figure 1.2.2: Plot of epitaxial film morphology vs. growth temperature and film composition for 1600 A thick GexSiI, films. (Bean et al., 1984)

1.3 Chemical Vapour Deposition

Chemical vapour deposition is among the most widely used film growth methods,

utilized in the deposition of critical device layers to the formation of lightweight carbon

fiber composite disk brake rotors for aircraft use (Meyerson, 1992). Characteristics such

as excellent film unifomiity, c o n f o d t y , mmpatibility with large area processing, and

relatively low apparatus costs compared with physical deposition techniques makes CVD

very attractive in the commercial market.

A low temperature process developed at IBM cded ultrahïgh vacuum chemical

vapour deposition (UHVCVD) will be discussed next. A schematic diagram of the

process is shown in figure 1.3.1. The UHV-CVD system is purnped at al1 times using a

turbomolecular pump, both when i d h g at base pressure. in the range of 1-5 x 1 oe9 torr,

and during film growth at 1 mtorr. The vacuum operation and high-density coaxial M e r

packing as in conventionai low pressure chemical vapour deposition (LPCVD), were

driven by the need of high-precision film growth. If films are to be accurate to several

Angstmms in thickness, the film growth rates need to be reduced to 1-10 Almin, and the

multi-wafer geometry of LPCVD is optimum to compensate for this (Meyerson, 1992).

Treatment of the silicon substrate before deposition is perfomed, is a requirement for

forming heterostructures of hi& quaiity. Rather than utilizing high-temperature oxide

desorption for surface preparation, UHV-CVD uses more conventional means. UHV-

CVD relies upon hydrogen passivation. An adlayer of hydrogen is formed in the process

of etching Si in hydrofluonc acid (HF), and was s h o w to reduce by 13 orders of

magnitude the reactivity of the silicon surface with respect to oxidants such as water and

oxygen (Meyerson et al., 1990). Therefore, in-situ cleaning methods as in MBE were

eliminated for UHV-CVD. The overall initial suface preparation for UHV-CVD is

straightforward involvïng a peroxide based chemical oxidation, followed by a 10:1 de-

ionized water/HF dip for 10 seconds. This pre-clean strategy produces a robust initiai

growth interface. As long as the hydrogen passivation is maintainecl, the silicon surface

remains free of contaminants.

Figure 1.3.1 : Schernatic of UHV-CVD process (Me yerson, 1992)

Afier the surface of the silicon substrate has been prepared for 10-35 wafers, held in a

quartz wafer boat, they are inserted into a load charnber which is pumped below 1 x 10"

torr. The wafer boat is then transferred under a hydrogen flow into the reactor, nominally

set to 550°C (- 500°C wafer temperature) at a pressure in excess of 200 rntorr. Al1

pumps remain in operation, resulting in a division of flow both down the reactor tube and

into the load charnber. This prevents the cross-contamination of the UHV side of the

apparatus by residuals in the load chamber. The entire cleaning and loading cycle for a

typical batch takes less than 15 min (Lafontaine et al., 19%). Upon isolating the growth

side of the apparatus, hydrogen flow is terminated, and silane flow is begun. At the flows

employed, typically 30 sccm (standard cubic centimeters per minute), system pressun

remains at - 1 .O mtorr.

The gaseous sources employed for the purpose of alloying or doping are germane,

diborane, acd phosphine, where these sources are diluted to produce films of deskd

chemical content. Helium is used as the carrier gas to facilitate diagnostics of the tlow

control systems, where it is straightfoxward to detect the He diluent gas compared wiîh

the highly diluent species. Film dopant content is linear in the atomic fiaction present in

the growth source, the same behg true for germanium incorporation (Meyerson et al.,

1987). Extremely precise dopant and d o y control is the result.

Auger electron spectroscopy (AES) and x-ray dihction (XRD) are used to measure

germanium concentrations. Photoluminescence (PL) is also used for M e r confirmation

of accurate Ge fraction. Thicknesses of deposited layers are determined by transmission

electron microscopy (TEM) and XRD. Background contamination effects and amount of

dopants incorporated are typically measured by secondary ion m a s spectroscopy (SIMS)

(Lafontaine et al., 1996).

1.4 Ion Implantation

Ion implantation is a process in which an ion of choice is embedded in the structure of

a selected material. ln a typical linear ion beam implantation (LIBI) (Figure 1.4.1). a

large voltage is used to bring the ion source niaterial up to a specified energy. For any

implantation to occur, the ions must be able to travel a distance of several meters without

collisions. This requires a vacuum of at Ieast 10-~ torr (Giedd et al., 1996). Severai

different kinds of purnps are utilized to achieve this pressure. Once the vacuum bas been

obtained, the ion source is filled with a gas. This gas is the ion source material, and can

be of a variety of compositions. The gas and ion source chamber are placed at a high

potential relative to ground. It is important that the entire source is well iosulated,

because at high potentials, dangerously potent arcs are possible (Hadey, 1976). A

current is run through the arc gap, aiming gas into a plasma.

The ions can be extracted h m the plasma using a magnet or strong electric field.

These extracted ions move as a beam towards the sample (Grovenor, 1992). The ions are

next passed through an analyzing magnet to reject those ions that are not of the desired

element. The operator can tune the magnetic field to allow onIy the ions of a certain

mass to p a s This beam then passes through an aperture that ennires the homogeneity of

the beam. At these apertures, a beam stop is present. The beam stop usually consists of a

large graphite block that on be dropped in fiont of the beam to block its passage. It is

here that an arnrneter is used to monitor the bearn current. The beam, screened of al1

undesired materials, is finally allowed to impact on the sample. The ions penetrate into

the material to some depth, altering the surface and bulk structure and the electncal

propenies of the matenal.

This method of doping materials allows for more exact control of where impurities

are placed in a material as well as more precise control of the impurity concentration in

the target material. Masks can be used to isolate areas of the target for doping. Lateral

motion of the dopants is much smaller in ion implantation since the diffraction of ions is

negligible (Giedd et al., 1996). The careful control of the beam current allows one to

monitor the dopant lrvels. Measurements of how many ions are impacting per unit time

c m be made. By multiplying the desired fluence by the entire scan area of the beam and

dividing by the bearn current, it is possible to calculate the implantation time necessary to

achieve the desired fluence.

The changes caused by ion implantation are ofien difficult if not impossible to

produce by other methods. The ability to control dopant parameters of the process aliow

the user to tailor the expairnent to produce correct concentrations and location of the

dopant materials.

Figure 1.4.1 : Photo of typical Ion beam implantation device (courtesy of University of Western Ontario)

1.5 Positron Annihilation S~ectrosco~v

The positron (e") which is an antiparticle of the electron (e'), has the same mass

within current experimental limits (51 1.003410.0014 kev/c2) and the same spin (%), but

with opposite charge and magnetic moment (Schultz et ai-, 198 8). It is stable in vacuum,

but rapidly themalizes in metais, and annihilates with an electron predominantly via 2-7-

ray decay (- 5 1 1 keV) with a mean lifetime that is typically only a few hundred

picoseconds (psec).

The positron's reaction in solids make it a very usehl method of studying defects on

the surface and in the bulk of materials When a positron from some radioactive source

enters a solid, it rapidly loses its kinetic energy (- 10 psec) until it is near thermal energy,

scattering between Bloch States to d i f i e through the solid. Figure 1.5.1 shows the

schematic of this process. The thermalbation tirne is short relative to the average

lifetime of the positron in the solid. Mer themalization, the positron rernains essentially

as a "free" or delocalized particle, although strongly correlated with conduction electrons

in its environment, until it annihilates in the bulk solid (Schultz et al., 1988).

DIFFUSION

POTENTIAL

Figure 1 .S. 1 : Schematic of positron's reaction in solids

The positron annihilation procedure is used to study lattice defects because a k l y

d i fh ing positron can localize in regions of minimum potential in a periodic lattice such

as vacancies as seen in Figure 1.5.1. This process is referred to as the Doppler

broadening of the energy of the annihilation y rays and relates to the positron trapping

into a vacancy. Figure 1.5.2 shows the effects of Doppler broadening (a) when a positron

is trapped in a vacancy, which also leads to longer positron lifetimes (b). Figure 1.5.2 (c)

shows how trapped positrons at defects, or at the surface can be parameterized by the line

shape parameter "Y, shown in the figure.

Energy

Figure 1.5.2: (a) Doppler Broadening (b) Positron lifetimes and (c) Formation of the S-parameter

The S-parameter is sensitive to vacancy concentrations of - 10-'crn-~. For each

implantation energy, the resulting S-parameter is

S = FbSb + FsSs + FdSd

where Fb Fs, and Fd conespond to the fiaction annihilating in the bufk, d a c e and at

defects and Sc, S,, and Sd are the S-parameters of the bulk, d a c e , and defect

respectively.

If a positron is trapped at a vacancy, it wiii be more likely to annihilate with the iow

momentum valence e1ectmn.s than the higher momenhim core electrons. Therefore the y

ray from the annihilation of trapped positrons will have a smaller Doppler shift and a

higher S-parameter.

The mean depth (D) at which positrons annihilate depends on their implantation

energy (E in keV). The group of Simpson et al. (1991) have simplified this relationçhip

as follows

D(in A) = 172 (E)'"

Positron beams al1 start with primary sources that have continuous energy

distributions arranged near a positron moderator or converter (Schultz et al., 1988). The

most cornmonly used primary sources are radioactive isotopes. A back-reflection

geometry (Figure 1.5.3) is typical with a '*CO source in front of a single-crystai

moderator to produce the bearn. Once the positron beam is extracted fiom the moderator,

it is magnetically-guided through a vacuum system to the target For the purposes of this

study, the positron beam utilized has a large Ge detector present for spectroscopy of

annihilation y rays for defect studies.

e+ Figure 1.5.3: Schematic of back reflection

geometry in e' beam production

1.6 Microsco~v Techniaues

Three different microscopy techniques have been utilized throughout this study of

defects in low misfit strained layer heterostructures. Each microscope possesses unique

characteristics which make them valuable tools for research These characteristics wiil

be discussed in detail in the following sections for (a) Nomarski Interference

Microscopy (NIM), (b) Field Ernission Scanning Electron Microscopy (FE-SEM), and (c)

Transmission Electron Microscopy (TEM). These microscopes are al1 located at the

Department of Metallurgy and Materials Science, University of Toronto.

1.6.1 Nomarski Interference Microscope (MM)

With slight modifications, a common optical microscope can be used for Nomarski

interference microscopy (NIM)(Figure 1 -6.1). A Wollaston prism and polarking lenses

are used to perform this procedure. Variations in surface level c m be seen by the

interference fringe contours or by changes in contrast performed by interference

microscopical techniques (Zhao, 1993). Coherent beams of light waves that are out of

phase by half a wavelength interfere destnictively which produce slightly displaced

images. These images cancel each other out at al1 points except those where the

displaced images are out of step, thus revealing the structure.

A parallel-sided double quartz pnsm .(Wollaston prism) produces the double image.

The angle of the wedges in the prism detemines the separation of the double image. An .

analyser which rotates in the column of the microscope produces coloured images

revealing certain wavelengths. This procedure illuminates and reveals desired structures

on the surface of the sample and can also provide optimum conditions for

photomicrographs by enhancing contrast.

I eyepiece

I I analyser

Inclincd giass slip illuminaior

objective lem sample

Wollaston prism

Figure 1.6.1 : Schematic diagram of Nomarski Interference Microscope (Haynes, 1 984)

1.6.2 Field Emission Scanninpr Electron Microscope (FE-SEM)

The Full capabilities of the FE-SEM will not be discussed here, however, certain

aspects of this powemil tool will be mentioned and their importance in this snidy. The

electron source in the FE-SEM is different fkom other conventional SEMs. Mead of

relying on high temperature to enable a fiaction of the free electrons in the cathode

material to overcome the barrier of the work function and leave (thennionic emission),

the FE-SEM utilizes an electric field at the tip to overcome the barrier (Goldstein et al.,

1981). The emitter for FE-SEM most commonly used is the <310> single crystal of

Tungsten (W). The 0 10> orientation has been chosen because the electmns were fomd

to travel in this direction most efficiently. Due to the Iarge mechanical stresses brought

on by the large electnc field at the tip (> 107 Vlcm), ody very strong materials can

withstand this stress without failing. Figure 1.62 displays the gun configuration of the

FE-SEM.

field emission tip 4 fmt anode a //\

Figure 1.6.2: FE-SEM gun configuration (Goldstein et al., L 98 1)

The three factors which govem the quality of a SEM image are: (a) electron probe

size, (b) probe brightness, and (c) interaction volume. A high intensity signai and small

probe size combine to produce the best imaging results.

Of the primary electrons that hit the sarnple fiom the incident beam, the two main

types of electrons which are most commoniy used for imaging are backscattered electrons

(BSE) and secondary electrons (SE). niirty percent of the pnmary electrons generate

BSE on average. BSE resdt fiom a shgie scattering event with the electron path greater

than 90' kom the incident beam direction. However, for the purposes of this study, only

SE will be outlined in detail. SE are the electrons most ofien used for sdace imaging in

the SEM. SE are defined as those electmns emitted h m the sample with an energy less

than 50eV (Goldstein et al., 198 1 ). interactions between the energetic beam and weakly

bound conduction electmns produce SE.

Due to the low kinetic energy of emission in which SE are generated, SE have a

shailow sarnpling depth ('; lm). SE are strongiy attenuated during motion in a solid by

energy loss due to inelastic scattering, which has a high probability in low-energy

electrons (Goldstein el al., 1981). The probability of SE escaping decreases

exponentiall y with depth in solid according to the following relationship,

P a exp (-dl)

Where z is the depth below the surface and A. is the SE mean fiee path. h is

approximately O S - 1.5 nm for metds and 10-20 for insulators.

Low accelerating voltages (< 5 kV) produce small interaction volumes in the SEM.

At accelerating voltages near 1 kV, spatial resolution can get as good as 5-7 A. Higher

SE yields are the result of lower accelerating voltages. Figure 1.6.3 shows the

relationship between the total emitted electron coefficient (backscanered and secondary)

and the beam energy. The secondary electron coefficient (6= nse/ne), where n s ~ is the

number of SE emitted fiom the sarnple and n~ is the number of beam electrons, rises to

reach unity at 1 kV. 6 increases to a value of 5 for nonmetais at 1-2 kV and then

continues to decrease to a beam energy of 0.1 at 20 kV for meîals.

The influence of specimen tilt on the SE cwfficient follows a secant law, 6(8) = 6.

sece, where 6, is S(B = 04 (Kanter, 1961). The SE CO-efficient increases with sec6. BSE

have a strong dependence of tilt angle. BSE increase with increasing tilt. Since SE are

also generated fiom the sample which suGers a backscattering event, the number of

secondary electrons will increase (Goldstein et al., 198 1).

LOG E,,

Figure 1 -6.3 : Relationship between total emitted e- coefficient and beam energy. (Goldstein et al., 198 1)

1.63 Transmission Electron Microscope CïEM)

Transmission electron microscopy (TEM) has proven to be a very powemil tool in

the study of defects in strained layer heterostructures. Individual layers can be isolateci

and observed with great accuracy in cross-section. Plan-view TEM allows one to scan

over large areas representative of the bulk matenal. Individual dislocations and

dislocation sources can dso be examined. There are situations where TEM results are

not representative of the bulk material due to structurai distortions induced by the shear

stresses which accompany any lanice parameter modulation. These distortions result

fiom TEM thinning during preparation in a direction perpendicular to the modulation

direction such that elastic relaxation of the internai stress-field responsible for the

tetragonal distortion occurs near the surface of the thin foi1 (Perovic, 1988).

For the study of dislocations and small defects such as precipitates or perturbations,

TEM is an indispensable tool. The combination of bright and dark field imaging

techniques and dif ict ion pattern adysis is essentiai in the chanicterization procedure.

Coherent elastic scattering produces difhction conhast DZhction contrast is simply a

special form of amplitude conh=ast beuiuse the scattering occurs at special (Bragg) angles

(Williams et al., 1 996). From the ciiffiaction pattern, which is usually known for single

crystal orientations, the 'NO-bearn' condition shouid be implemented. Tilting the foi1

and obtaining a recognizable symmetry is a convenient starting point (Thomas et al.,

1979). Because of the 180' ambiguity in spot patterns, the spot pattern by itself does not

yield a unique foi1 orientation. However, a Kikuchi pattern can be indexed uniquely and

is facilitated with a Kikuchi map. For vacancy or interstitiai loops, defects or faults, this

information is needed to d e t e d e the sense of the strain fields.

Defects with zero or small strain fields (e.g. voids, inclusions of second phase with

zero misfit) are visible only through "scattering factor" contrast-columns through such

defects appear to be of different thickness from those through adjacent regions of matrix

crystal. They are thus only visible in situations where strong thickness fringes are seen,

that is, in thin areas when the crystal is at the reflecting position (Thomas et al., 1979).

Defects with greater strain fields produce small "black-white" contrast (small dark

areas on a light background) when imaged at deviated positions. This black-white

contrast is oscillatory with the depth of the defect (with a penod equal to that of thickness

fnnges) (Thomas et al., 1979). The sense of the g-vector is useful in determining the

nature of the defect (interstitial or vacancy-type). Smdl Ge perturbations found in Ge-Si

strained layer heterostructures were determineci to be interstitial in nature due to the

contrast on the perturbation which went fiom light to dark following the g-vector in the

[220] direction (Perovic et al., 1992).

CHAPTER 2 - EXPERIMENTAL

2.1 Heterostmcture Growth Conditions

The two growth techniques used in this study were MBE and UHV-CVD. Both of

these facilities are located at the National Research Council of Canada ( M E C ) in

Ottawa. The heterostruchires were grown on 4 inch diameter (100) Czochralski Si

substrate wafers (6 inch for CVD). The Czochralski method is show in Figure 2.1.1.

Figure 2.1.1 : Czochralski growth method

UHV-CVD conditions for growth were taken fiom Lafontaine et al. (1996). The

UHV-CVD samples were grown using a "Sirius" hot wall UHV-CVD reactor. The

growth chamber consists of a quartz tube heated by a hirnace and evacuated by a

turbmolecular/roots blowedro tary pump system. The base pressure was 1.5 x 1 0" mbar

at T = 52S°C. Silane (100%) and germane (10% in helium), of Matheson ULSI grade,

were used as precursors. Customdesigned mass flow controlles (MKS, type 1449A)

were used to control gas flows.

M e r a standard "RCA" clean and an HF:H20 (1 : 10) dip for 10 seconds, d e r s were

introduced in the chamber through a loadiock which was pumped d o m to a base pressure

of 1 O" mbar in 5 minutes. A hydrogen flow of 300 (sccm) was used during d e r

transfer into the main chamber in order to prevent contamination from the loadlock

Growth was initiated immediately after the t r d e r is completed, Si& injection. The

entire cleaning and loading cycle for a typical batch takes less than 15 minutes. Table

2.1.1 shows the growth parameters of the UHV-CVD samples used.

(see Appendix for definition of reK, N, h, and H)

The MBE procedures foliowed those of Baribeau et al. (1994). GeXSii.JSi

superlattices were produced in a VG Semicon V80 MBE system. Unlike UHV-CVD,

only one wafer was grown at a tirne in this system. A 4 inch wafer was placed face down

inside the ultra-high vacuum chamber. An in-situ pre-clean was perfomed on the wafer

above 850°C under 0.1 A/s Si flux for 15 minutes. The wafer was then heated to the

desired growth temperature (360-500°C) (see Table 2.1 . 1 (b) for details). Each wafer

received a 15 nrn buffer layer of Si prior to growth of the superlattice. Samples 1688-

1703 were grown at rates of 5 &sec. Samples 1706 and 1707 were grown at slightly

lower growth rates. The growth rates were controlled by a Sentinel III in situ monitor.

The vacuum pressure during growth was maintained at 5 x IO-'* torr. Accurate

TC f i

( M W 1 03 57

composition readings were made by doublecrystal x-ray rocking curves.

N Periods

10 10

Sample #

CVD-9 CVD-61

Si)

21 29

X

0.10 0.13

h(siGe)

(nm) 15 7

Table 2.1 .llbl- Structurai Panuneters for MBE Ge,Sil.&

2.2 Ion Implantation Conditions

Sample#

MBE1688 MBE 1689 MBE1690 MBE1691 MBE1702 MBE 1703 MBE 1706 MBE 1707

Poçt-growth ion implantation was performed on four of the ten wafes produced.

CVD-9 and CVD-61 were taken h m the UHV-CVD sarnples and MBE1702 and

( s e Appendix for definition of T = ~ , N, h, and Fi)

Growth Temperature

CC) 405 405 360

MBE 1703 h m the low temperature MBE samples. Ion implantation procedures follow

those of Goldberg et ai. (1993) and Labrie et al. (1996). The GexSii,/Si superlanices

x

0.9 0.9 0.9

were self-implanted with 540 keV Si ions to a fluence of 2 x 1014 ions/cm2 ushg a 1.7

MV Tandetron accelerator located in the Department of Physics, at the University of

h(='

(m)

15 10 15

360 450 500 400

Western Ontario. Wafen were secured to a nickel block with thermally conducting paint

0.9 0-85 0.84 O. 12

to ensure good thermal contact. The temperature of the block was maintained at 25 k 1°C

IO 15 17 11

during irradiation. An implant flux of - 1 15 &cm2 was kept constant throughout the

Gtr

@Pa)

103

SI)

(m)

21

IO 10 10 I O

30 23 24 15

bombardment. Channeling effects were avoided by rotating the sample - '7' about axes

N Periods

10

47 89 97 137

10 14 425

both perpendicdar and parallel to the incoming beam.

30 21

147 0.13 1 10

10 I O

47 1 03

23 Rapid Thermal AnneaIin~ (RTA)

To eliminate kinetic barriers, pst-growth annealing via RTA was performed on

CVD-9, CVD-6 1, MBE-1702, MBE-1703, MBE-1706, and MBE-I 707. Table 2.3.1

shows the RTA temperatures and h e s for these samples. These temperature/time

combinations were chosen according to the equivalent strain relaxation contours of

Houghton et d (1993). The RTA were performed using a Heatpulse 410, in a dynamic

nitrogen atmosphere. The Heatpulse 410 is located at NRCC, Ottawa, To avoid

contamination of the samples fiom dust particles, the RTA experiments were conducted

in controlled airflow cleanrooms and full body suits had to be wom, as they were for

CVD and MBE growths. Small 1 cm x 1 cm samples were placed on a 4 inch diameter

Si wafer substrate which was used as a stage in the Heatpulse 410. RTA temperature

ramping was controlled via cornputer software.

Table 2.3-1 - Temperature and T h e Parameters for RTA

2.4 Nomarski Interference Microscopy (NIM)

Misfit dislocation nucleation rates and densities were calcuiated Eom images

produced on an optical microscope, modified for Nomarski imaging. Specimen

preparation for NIM was performed on the samples used for RTA. The 1 cm x 1 cm

wafer pieces were immersed in a dilute Schimmel etch (4 parts 48% HF : 5 parts 0.3-M

CQ) for desired etching times. The etching times were calculated using a Dektak 3-30

Temperature CC) 700 800 850 900 950 1 O00

RTA time (s) 1 O0 100 30 10 5 5

Version 2.13 depth profiler. Small drops of a black wax, immune to the etching solution,

were placed in the corner of each wafer. After submersion in the Schimmel etch for

approximately 30 seconds, the samples were profiled in the Dektak across the

rtchedhon-crched boundary line. Etching rates were determined for each sample. The

misfit dislocations nucleated at the interface of the substrate and the first Ge-Si layer,

therefore the proper rtching depth to reveal the rnisfits was calculated by using the

geometry of the layes (e-g. for CVD-9: 10 x (21 nm Si + 15 nm GeSi) = 360 nrn total

height to be etched). Sarnples needed for TEM observation were not etched to preserve

the lattice planes.

Figure 2.4.1 shows a typical Nomarski micrograph. Care was taken to observe

representative areas (several cm') away From scratches or edges to determine the nurnber

of nzw misfit segments formed per unit area per unit time. Dislocations can nucleate

from scratches. clsaved edges and other stress concentrations and cause inaccurats

readings of the tme strain relaxation data of the heterostructure.

Figure 2.4.1 : Micrograph of a typical Nomarski image showing UHV-CVD sample RTA for 100 s at 700°C. Individual misfit segments can be seen.

2.5 Field Emission Scanning Electron Mieroscom WE-SEMI

FE-SEM exgnhtion was carried out on the S4500 FE-SEM. Specimen preparation

for the FE-SEM was the easiest to carry out compareci to al1 other techniques. MBE-

1688, MBE-1689, MBE-1690, and MBE-1691 were the samples used for FE-SEM

observation Small(2 mm x 3 mm) pieces were cleaved using a diamond tip cutter. No

other preparation procedures were required. The fkeshly cleaved edge was placed in the

rotating sample holder and loaded into the vacuum chamber of the FE-SEM.

Low accelerating voltage (1-2 kV) experîments were conducted on the samples with

4-6 mm working distances. The secondary electron (SE) detector was used to examine

the surfaces of the MBE growths. The usual stigrnator corrections were performed to

pmduce the high-resolution images. Slow scanning speeds were used to obtain the

highest level of detail and contrast. Micrographs were acquired using the Quartz-PCI

software present in the FE-SEM laboratory.

2.6 Transmission Electron Microscopv (TEM)

A Hitachi H-800 TEM operating at 200 kV was used for specirnen characterization.

Sarnples were prepared for both cross-section and plan-view anaiysis. The cross-section

procedures are as follows. Several strips (1 mm x 5 mm) were cut fiom the 400>

oriented wafen using an automated cross-cut saw, located at NRCC. The strips were h t

immersed in acetic acid and then ethanol to clean the samples of any contaminants. Next,

two strips were turned 90' so the fih sides faced each other. A silver epoxy was used to

glue them together. This produced <O 1 1> surface normals. See figure 2.6.1 for

schematic of procedure. The epoxy was then cured at 100°C for approximately 3 hours.

The next step was to core the sample using a Precision Instruments drill. The sample

was placed on a g l a s slide and kept there using a fast cooling wax A 3 mm diameter

disc was cored using a 1 pm particle diamond paste slurry. The disc was then mounted

ont0 a polishing tool using the same wax. The disc needed to be mechanically dimpled to

a thickness of - 50 pm i m g the 1 pm diamoed paste. A D500 Dimpler was used for this

procedure.

The final step for the cross-section specirnen preparation was ion-milling. A Gatan

Ion Mill was used for this procedure. Argon ion sputtering was performed on both sides

of the disc at the same time until the sarnple was electron-transparent in the region of

interest. A sputtering angle of 15' and a voltage of 4.5 keV were used for the first 5

hours of operation to mil1 through the buik of the matenal. A shailower angle of 1 1' was

then used to min& the depth of the amorphous surface layers which result nom

mechanical damage in the sputtering process (Perovic, 1988). If done correctly, the

cross-section sample had four areas of observation. The entire process can take up to 25

hours per sarnple Cor TEM preparation.

/ ' / / A

slice sections

1 mm size strips

SOO A thick Figure 2.6.1 : TEM cross-section

preparation

The plan-view sampies were easier to prepare because there was no epoxy present,

therefore no weak adhesive bonds to break. A 4 mm x 4 mm sample is placed face down

on a glass slide and glued with the same wax as above. The coring, dimpling. and ion-

milling procedures were al1 the same as above. The only clifference is that fih side of

the sample never cornes in contact with the dimpling wheel or the argon sputtering stream

of ions. Thus, oniy the backside of the of the plan-view sample was ion-rnilled at a tirne.

prolonging this procedure.

In order to produce quality TEM micrographs for observation, detailed procedures

had to be followed. After producing suitable sarnples for the TEM, the next step was to

perform TEM imaging. For the purposes of this study, acquiring a working difhction

pattern (DP) and obtaining the mo-beam condirion were essential for TEM observations

to be made. Figure 2.6.2 shows (a), a schematic of the DP of a diamond cubic lattice

with [O011 beam direction and (b), an actual image of the pattern obtained in the TEM

with g = [400]. The beam direction or zone-axis z = [OOI] was used for plan-view

images and z = [O 1 1 ] for cross-section images (Figure 2.6.2 (c)).

To obtain good strong diffraction contrast in both bright field (BF) and dark field

(DF), the specimen was tilted to two beam conditions. in which only one difhcted beam

was strong as seen in figure 2.6.2(b). After obtaining the two-bearn condition, the

objective aperture was centered on the incident b a r n to produce BF images and centered

on one d ihc ted beam to produce DF images of the desired g-vector orientations.

Strain field contrasting was also performed on the TEM. In order to locate small

perturbations in the stained layer superlattices (SLS), first discovered by Perovic et al.

(1990). lattice stra in effects around the perturbations had to be examuied These lobes of

low intensity are symbolic of strain fields (Williams et al., 1996). Strain field contrast

images were obtained in DF mode.

Figure 2 - 6 2 (a) schematic of DP of diamond cubic lanice (beam direction, z = [O0 11) (b) actual TEM micrograph with g = [400] displaying 2-beam condition (c) schematic of DP showing cross-section arrangement (z = [O 1 11)

2.7 Positron Annihilation Conditions

The positron laboratory is located at the University of Western Ontario. It was here

that al1 the S-parameters were modelled and statistically calculated. One of the two

variable-energy positron beams was used to characterize the defect structures in the

samples (i.e. voids, vacancies, etc.). The positron beam started at the primary source,

which had a continuous energy distribution arranged at the positron moderator. The

radioactive isotope used was the S 8 ~ o source in front of the single-crystal moderator.

Magnets existing in the apparatus guided the beam through the vacuum chamber to the

target A large Ge detector was used for spectroscopy of the annihilation y rays for defect

c haracterization.

The S-parameter, discussed in Section 1.5, was determined for each sample in which

positron annihilation was performed. The correspondhg equations were also determined

(i.e. the mean depth at which positrons annihilate). By using these equations, different

positron annihilation models were calculated and defect concentrations were determined.

CHAPTER 3 - RESULTS AND DISCUSSION

The following section considers the observations and results obtained h m the

experiments and characterization techniques. In this examination of seain relaxation

effects of low misfit heterost~ictures, evidence of dislocation multiplication has not been

obtained. The experimental procedures outlined thus far have centered around seain

relaxation effects in the fim stages of relaxation (i.e. Stage 1 regime, dislocation densities

between O - 10' cnf2). It is in this area that there is little known about the precise

injection mechanisms of misfit dislocations. Multiplication processes and dislocation

interaction effîcts can be ignored in this regime. It is not until dislocation densities > 10'

that these issues are of concem. Dislocation interactions introduce a host of new

problems (Le. dislocation blocking, pile-ups, etc.) in trying to detemiine s t n h relaxation

mechanisms (Schwarz, 1997, Schwarz and Teaoff, 1996).

A novel method for controlling the generation of misfit dislocations by point defect

injection via ion-implantation in m e d Iayen has been examined. Bulk measurements

were made to detennine the misfit disiocation nucieation rates and densities for various

Ge,Si&i (x 5 0.13) heterostructures grown by MBE and W - C V D methods.

Nomarski interference microscopy was used for these bulk measurements. TEM has

been used to locate sources of nucleation and to examine the dislocation structures both

in plan-view and in cross-section.

Positron annihilation was perfomed on samples which were thought to have

substantid vacancy concentrations. On some of the low-temperature MBE growths,

where incoherent growth resulted, positron annihilation was used to mode1 the void

distribution in the niperlattices with the S-parameter. Along with FE-SEM. the void

study of strained layer superlattices (SLS) yielded some interesting resuits which were

consistent with a "morphological instability phase diagram" proposed by Perovic et al.

(1 993).

3.1 Strained Laver Geometrv

The strained layer geomehies play a major role in detennining how the

heterostructure will relax and through what mechanisrns. Figure 3.1.1 illustrates the

typical geometrical distribution in the SLSs studied. Each individual strained layer (i.e. h

or H for GeSi and Si thicknesses, respectively) was grown below their criticai thickness.

N represents the nurnber of periods of repeating (h + H) layers. According to this

configuration, compressive strain is generated in the layers as growth proceeded and

relaxation was attributed to the extension of misfit dislocations at the first GexSil,,/Si

interface. The unbalanced force responsible for the misfit dislocation nucleation and

propagation in strained epitaxial layers exceeding equilibnum critical thickness is the

effective stress, r,n (see Appendix), which was k t defmed by Matthews, Mader, and

Light (Houghton, 1991). This effective stress represents the driving force for misfit

nucleation and propagation.

A range of reff was examined in this study to obtain varyhg nucleation rates and

misfit densities. The effective stress is detemiined by the imbalance between the

resolved shear stress acting on the threading dislocation slip system and the iine tension

in the extending a/2 < 1 1 O> misfit dislocation segment (Houghton, 1 99 1). Typically,

SLSs grown with higher r,a values displayed greater nucleation rate and misfit

dislocation density values. However other factors were aiso important in detamining the

nucieation rates and dislocation densities.

Figure 3.1.1 : Geometrical distribution in the superlattice

3.2 Bulk Ouantitative Measurements

Large area diagnostic techniques were used to determine the misfit dislocation

nucleation rates and mis fit dislocation densities of the heterostructures. A comprehensive

strain relaxation rnodel (Houghton (1 99 1 ), Perovic and Houghton (1 993)) has been

developed based on bulk measurements of rnisfit dislocation nucleation and glide for a

range of metastable (1 00)aiented Ge,&,/Si (0.03-0.25) heterostructures grown by

MBE, RTCVD, and UHV-CVD. The same tests were conducted in this study following

their mode1 using GexSii../Si (0 .09~0 .13) heterostructures grown by MBE and UHV-

CVD.

Post-growth annealing (5-100 s, 700-1000°C) was carried out on the metastable,

coherently suained superlatticw (i.e. initially misfit dislocation-free) according to the

procedures outlined in Section 2.3. Afier a Schimrnel etch, quantitative measurements

were conducted on the samples (severai cm2) away fiom edges or other stress

concentrations. Therefore, the number of new rnisfit segments generated per unit area

Figure 3 2 . 1 : Nomarski images of initial stages of relaxation for CVD-6 1 after RTA for (a) 100 s at 700°C (b) 100 s at 800°C ( c ) 10 s at 900°C and (d) 5 s at 1 OOO°C

per unit t h e were determinecf for the nucleation rates (dN(t)/)/dt) (sez AppAppe in (cm-2

s-') for a given driving force (snr). In order to calculate the misfit dislocation densities

(N(f)) in (cm-'), the inverse of the distances between the misfit dislocation lines for a

given rnagnification were measured.

These values were measured directly off the Nomarski micmgraphs such as in Figure

3 -2.1. This figure illustrates the evolution of orthogonal a/2 cl 1 O> misfit dislocations for

the CVD-6 1 multilayer of Table 2.1.l(a). At 700°C, little misfit injection has o c c d

whereas at 800°C a low density of misfit segments is apparent The two higher anneals

(Le. 900°C and 1000°C) indicate a rapid increase in both average misfit segment length

and an increase in the density of activated nuclei at elevated annealing temperatures.

Figure 3.2.2 displays the nucleation rate vs. inverse temperature (KI) for various

sarnples grown by UHV-CM and MBE processes. An important fact to consider here is

that although the nucleation rates Vary between samples, due to layer geometry and

effective stresses, the activation energy for nucleation is consistent at a value of Q. = 2.5

+ 0.5 eV. This value is found by taking the dope of the line on the nucleation rate curve.

It represents the exponentiai evolution of misfit dislocation density (N(t)) as a h c t i o n of

temperature (7). This value is consistent with the work done by Houghton (1991) and

Zhao (1993) and recently by the group of Wickenhêuser et al. (1997). n i e group of

Wickenhauser et al. (1 997) were able to isolate the three different mechanisms involved

in stmùi relaxation (Le. nucleation, propagation, and multiplication) and study them

independently by selective epitaxy. The group of Tanaka et al. (1996) found the

activation energy to be 2.5 eV by studying photoluminescence (PL) spectra of deformeci

bulk Si-Ge alloys. By studying the change in alloy compositions around dislocations,

caused by elastic interactions between constituent atoms (i.e. Si and Ge), pecuhar peak

shifts of D 1 and D2 lines were observeci.

Ternperature (OC)

7.6 8.1 8.6 9.1 9.6

Ternperature (1 /K) (1 E-4)

Figure 5 -2.2: Nucleation rates of misfit dislocations vs inverse temperature

Other attempts have also been made in trying to determine the strain relaxation

mechanisrns. The group of Hull et al. (1989) used in situ TEM measurements to produce

activation energies for strain relaxation. Their low values for activation (Le. Qv = 1.1 eV

- 2.2 eV and Q,, = 0.3 eV for glide and nucleation, respectively) indicate that other

mec hanisms are taking place (i .e. dislocation interactions). Due to the srnall dimensions

of the thin foils used in the TEM, oxidation-induced vacancy injection at the surface cm

occur and vacancy migration in the Si substrate wodd be the rate-limiting process and

would account for the Qn = 0.3 eV measured by Hull et al. (Perovic et al., 1995).

It is also important to note the nucleation rate incrrases with t h u g h a power-law

dependence, n = 2.5. The nucleation rate varies hearly with the initial misfit dislocation

source density (No) present in the as-grown material at t = O (Le. subsîrate threading

dislocations, residual substrate-baer layer precipitates, etc.). Different substrate

cleaning procedures pnor to growth result in varying values of the initial source density,

No, for MBE (No > 10) cnf2), and UHV-CVD (No S 102 cm"). The linear increase with

time proceeds to densities well beyond No indicating the themial activation of new

sources in the early stages.

Figure 3.2.3 displays a plot of misfit dislocation density vs. inverse temperature for

UHV-CVD and MBE samples. In these samples the activation energies were found to be

Qr = 4.2 + 0.5 eV. This value is double that of misfit glide and close to the value of 4.75

eV reporteci by Houghton (199 1). The reason for the higher activation energy is that this

value takes into account the overdl strain relaxation in this stage of relaxation. Therefore

the nucleation and glide energies are in series to produce this value (i.e. Q, = Q, + Q,).

Figures 3 2.2 and 3.2.3 represent measwements taken from the initial stages of strain

relaxation (i-e. Stage 1 regime). This regime is said to be nucleation-limjted, b u s e the

misfit dislocation nucleation increases linearly with tirne, beyond No. Houghton et al.

( 1 990) found that the misfit dislocation velocity was effectively constant which resdted

in a linear increase in the ovemll seain relaxation rate with time. The rate-limiting step

in rnisfit dislocation nucleation has k e n attributed to vacancy formation andor migration

during the generation of the incipient dislocation loop at growth-induced interstitial

perturbations (i.e. Ge-rich platelets) which will be discussed in detail in Section 3.3.

Temperature ("C)

Figure 3 -2.3: Misfit dislocation density vs. inverse temperature for UHV-CVD and MBE matenal

Controlling misfit dislocation generation in strained layer superlattices has been a

major concern for researchers and manufacturers of electronic devices. To achieve

optimum performance fiom high-speed GeXSiiJSi strain layer devices, the retention of

structural perfection is a necessity. A novet procedure for controlling and reducing misfit

dislocations in these structures has been examined here. A detailed shidy on G+Sii,/Si

as-grown heterostructures verses 540 keV self-irradiated Si samples has been conducted.

Figure 3.2.4 shows a plot of misfit dislocation nucleation rate data acquired for the

UHV-CVD matenal shown in figure 3.2.2 (Le. CVD-61) and a plot of the same material

implanted with 2 x 1 014 Si ions/cm2. Figure 3 . î .S shows a Nomarski micrograph of the

CVD-6 I heterostnicture. An intentionally introduced scribed line, to assist nucleation of

new misfits, separates the as-grown and implanted sides. After a RTA of 850°C for 30

seconds, the following features were observed. Note how the misfit dislocations extend

Temperature (OC)

7.6 8.1 8.6 9.1 9.6

Temperature (1 /K) (1 E-4)

Figure 32.4: Nucleation rate data for UHV-CVD as-grown and implanted material

Figure 3 -2.5: Nomarski micrograph showing as-grown and implanted sides of UHV-CVD matend with rnisfits nucleating from scribed line. The scribed line separates the implanted and as-grown sides (see text)

M e r into sample on the a s - p k side. Wiîh time and increased temperatUres, the

misfit lines will extend across the entire sample. The scribeci line was introduced after

implantation to separate the two sides and RTA was then subsequently performed.

Figure 3.2.6 shows a plot of the misnt dislocation density vs. inverse temperature.

The dislocation densities dif5er by as much as a factor of eight for CVD-61 and a factor

of five for CVD-9 (not shown here)- A dramatic decrease has been observed in the

nucleation and densities of misfit segments on the implanted sample. However, the

activation energies for nucleation and overd strain relaxation remain constant at Qn = 2.5

+ 0.5 eV and Q, = 4.2 t 0.5 eV respectively, uidicating that the same mechanisms are

present even afier implantation.

Temperature (OC) 1000 950 900 850 800

7.6 8.1 8.6 9.1 9.6

Temperature (1 /K) (1 E-4)

Figure 3.2.6: Misfit dislocation density data for CVD-6 1 (as-grown and implanted)

The results shown here are consistent with a mechanism where Ge-rich platelets,

which possess an interstitiai eIastic field, can be relieved by the injection of point defects

and hence reduce the overall stress concentrations surrounding the perturbation which

assist the rnisfit dislocation generation Therefore, an overd decrease in sûaîn relaxation

via misfit dislocation is expected, which is in agreement with the resuits.

There are many theoretical and experimental atomic diaision studies about Si and Ge

and well-defined values for energies of vacancy formation (Er) and migration (E,). Van

Vecheten (1980) observed values for E~'' = 0.33 eV, E:' = 2.3-2.5 eV, E? = 1.0 eV,

and E? = 2.0 eV. Therefore, with the Q, = 2.5 f 0.5 eV found in this work, the energy

term can be attributed to: (a) vacancy formation in Si or (b) vacancy formation in Ge

followed by migration in Si. The latter is consistent with these resdts. This information

lead to the next set of experiments.

It is now accepted that the e l d c stress-field smunding the interstitial platelets (i.e.

Ge perturbations) can be relaxed by condensation of vacancies, which would Ieave

behind a vacancy-type pnsmatic loop. Section 3.3 will show evidence of the Ge-platelets

and the resultant behavîour after RTA-

With a large enough vacancy supersaturation during growth of the heterostructure,

dislocation loop nucleation will be controlled solely by vacancy migration. Vacancy

formation (E:' = 2.0 eV) wouid not be required. Therefore, the activation energy for

nucleation (QJ wouid be sufncientiy reduced near the value of E~'' = 0.33 documented

by Van Vecheten (1980). Low temperature MBE growth procedures would have to be

utilized in order for enough vacancies to be grown into the materid. However, if the

growth temperature is to low and the growth rates are not modifieci, nano-scale void

formation will r e d t (Section 3.5) due to v m c y condensation. In order to produce

these structures with a sufnciently large vacancy concentration, the growth rates of Si and

Ge would have to be decreased to d o w enough tirne for the surface migration of atoms

on the substrate at these reduced growth temperatures (Le. 400-450°C). UHV-CVD

couid not be used for these low temperature experiments because the temperature mut

remain constant at approximately 52S°C to maintain a base pressure of 1.5 x 10" mbar

necessary in the growth chamber.

Figure 32.7 shows the misfit dislocation nucleation rates for the CVD-61 material

and a low-temperature MBE material (i:e. MBE-1707). It is obvious here that a different

stmin relaxation mechanism is fiinctioning in the low-temperature MBE matenal. The

slope of the plot of this material was reduced giving an activation energy of 0.5 f 0.05

eV. Therefore, a suficiently large vacancy supersaturation h a . dominated here with

respect to the other structures grown at higher temperatures and growth rates. Vacancy

migration is the operative mechanism for low-T MBE material. The densities of misfit

dislocations were also affected as seen in Figure 3.2.8. The resuiting activation energy

dropped to 2.0 f 0.5 eV, exactly 2.3 eV lower than the overall strain relaxation values for

UHV-CVD matenal. Thus, for the low-T MBE growths, 0.5 t 0.05 eV c m be attributed

to the vacancy migration. These values are consistent with the predicted strain relaxation

mode1 proposed above.

To get a more accurate assessrnent of the functioning mechanism, positron

annihilation spectroscopy has ken conducted on the MBE material and the as-grown

CVD-9 matend as a control specimen, due to the low density of initial grown-in defects

Temperature (OC)

7.6 8.1 8.6 9.1 9.6

Temperature (1 /K) (1 €4)

Figure 3 -2.7: Nucleation rate data for UHV-CVD and low-T MBE material

Temperature (OC)

7.6 8.1 8.6 9.1 9.6

Temperature (1 /K) (1 €4)

Figure 3.2.8: Misfit dislocation deasity data for UHV-CVD and low-T MBE material

Depth (angstroms)

Positron energy (keV)

Figure 3.2.9: Positron Annihilation Spectmscopy data for low-T and UHV-CVD material

(No 1o2 cm'2). According to Figure 3.2.9, the S-parameter has been modeled for the

various MBE structures. The S-parameter in the graph corresponds to the ratio of the

number o f y-rays in the central region of the peak (5 10.27 to 5 1 1.73 keV) to the total

number in the peak (504.51 to 517.49 keV). The lower the Doppler bmadening, the

higher the S-parameter-

Ail the samples have the same S-panuneter for high positron implantation energies

indicating that there are no signifi~ant ciifferences deep in the samples. For some of the

samples (1702, 1703, and 1706), the graph is less steep in the beginning for positron

energies less than 2 keV. This means that the hction annihilating at the sdace does not

change much for increased implantation energy. This codd be caused by an electric field

driving the positrons back to the surface. Another possible expianation is that the= codd

be a layer at the surface which is different in some way, like a layer of oxide. This codd

be around 400 A for sample 1702. CVD-9 could be fit well with a model having no

defects. However, 1707 could be fitted using a model with a different sudace S-

parameter and assiiming a fiactional vacancy concentration of about 5 x 1017

(Simpson et a[., 1 994). Therefore, MBE- 1707 may have displayed a behaviour caused by

a high vacancy concentration after growth, which is consistent with the bulk

rneasurements conducted (Figures 3.2.7 and 3.2.8).

3.3 Sources of Nucleation

The sources of nucleation and the subsequent expansion of misfit dislocations have

been studied by many researchea. Perovic et al. (1989) found that coherent /?-Sic

precipitate plates, localized at the Ge&,/Si interface acted as efficient sources of 60'

misfit dislocations. Hurnphreys et al. (1991) observed the "diamond defea", which was

a faulted disiocation loop which acted as a heterogeneous source of 60' misfits.

However, heterogeneous sources like these can be controlled and are a result of specific

materials systems or growth conditions. Therefore, a generaüzed nucleation process for

low misfit heterostructures was required that was consistent with measined activation

energies regardless of growth conditions and procedm (Le. MBE CVD, RTCVD etc.)

for low misfit heterostructura.

The formation of dislocation loops in the interface has been recognked as the widely

accepted source for nucleation of the misnt segments (Jhmat et al., 1990, Jain et al.,

1995, Zou et al., 1996). Hirsch (1997) examined the conditions under which the Ge-

platelets act as sites for nucleation of "double half-loops" which expand in the epitaxial

layer and act as sources for the 60' misfit dislocations. Hirsch's theory of nucleation is

consistent with the mode1 developed by Perovic and Houghton (1992) and has dso been

examined here.

The sources of nucleation for the dislocation loops and the expansion of the incipient

loop to fonn stable misfits are attributed to the loss of coherency at interstitiai

perturbations (Le. Ge-rich platelets). Figure 3.3.1 shows a TEM image obtained in strain

field contrast of the Ge-nch perturbations. The g = 12201 was used to find the small

platelets (diameters - 1.5 nm, - 2-3 monolayers thick). DF imaging was used for the

strain field images because of the anomalous absorption effects. The high intensity

region corresponds to a region of good transmission where Bloch wave II predominates

since its intensity is concentraied between the difbcting planes. The adjacent region

trmsmits electrons poorly since Bloch wave 1, with its intensi~ peak at the dificting

planes, dominates and thus appears dark relative to the background (Perovic et al., 199 1 ) .

Images in BF are symmeîrical and therefore, it is more difficult to show this wntrast

behaviour. The contrast over the platelet goes h m Light to dark in the direction of the g-

vector, which means that the platelet is interstitially strained. Plateiet formation was

assumed to be an inherent feature of MBE growth (Rowell et ai., 1993). It was assumed

that the hydrogen-passivated alloy surface typical in CVD prohibited these features From

forming. However, platelets were indeed found in the CVD material as well as the MBE

material. The suain relaxation mode1 held for both growth techniques as seen in Section

3.2.

The group of Rowell et al. ( 1993) found that upon tilting the sample in the TEM to

othsr zone avrs other than the [001], the associated image conuast relationship indicated

that the perturbations are atornic-scale regions possessing plate-shaped symmetry and a

Figure 3.3.1: TEM dark field image of Ge-rich platelets in CVD-9 sample (g = [220])

larger larrice parameter than the adjacent matrix region. They found perturbations in a

wide range of single and multiple layer heterostructures. The defects were never

observed in the thin foi1 regions away h m the heterostnicture thus d i n g out specimen

damage-related artifacts as possible defect sources. A broad photoluminescence (PL)

peak, shifted lower in energy than the phonon-resolved PL, has been assigned to the

formation of these perturbations indicating an increase in the strain energy density. The

shift occurs since excitons in the Ge&, are localized in the lower band-gap Ge-rich

platelets (Rowell et al., 1993).

With the use of an energy balance approach (see Appendix), Perovic and Houghton

(1992) determineci that a criticai radius condition can be detennined where the strained

perturbations can iower its energy upon creating a prismatic loop at its interface. It was

also found that for large misfit strains (Le. x 2 0.85, for GexSii,/Si), Ge platelets

embedded in Si will homogeneously generate an interfacial dislocation loop (vacancy-

type) without having to overcome an activation energy through pst-growth ameaiing.

Hence, the term "banierless" misfit dislocation nucleation was coined- At Lower rnisfit

strains, the critical radius is important, and requires pst-growth annealing to generate

misfit dislocations at the interface. Hirsch (1997) has also argued that loop nucleation is

more likely in larger perturbations and that the rate controlling process is the diffusion of

vacancies to the perturbations.

After the generation of the prismatic dislocation loop, the nucleation of the half-loops

and their expansion in the superlattice is the next step (see Figure 1.1 4. In accordance

with Hirth and Lothe (1 982), it is possible that prismatic vacancy loops on ( 100) planes

can act as sources of dislocation shear loops on (1 11) planes in the presence of a

resolved stress on the glide plane (i.e. retr). Therefore, segments of the { 100) prismatic

loops can bow out ont0 the { 1 1 1 } planes, and form haif-loops.

Perovic and Houghton (1 992) used the total Helmoltz energy equation: E, = E, - Ed t

Es - TS, where Ed is the dislocation half-loop self-energy, E, is the misfit strain energy, Es

is the surface step energy, and TS is the entropy of the loop (see Appendix for the

complete goveming equations). The activation energy for nucleation of misfit

dislocations is obtained from this expression by substituting in the cntical radius (R') of

the Ge perturbation. Therefore there will exist a range of effective stresses, increasing

with perturbation size that facilitate a driving force for reduction of the activation barrier

for rnistit dislocation injection. [t is these sources that f o m misfit dislocations after

RTA. This mode1 is consistent with the results found in this study.

Figure 3 - 3 2 and figure 3 - 3 3 are TEM images of CVD-9 after RTA for 5 seconds at

1 OOOUC. Figure 3.3.2 is a BF image of orthogonal misfit expanding out from an original

nucleating source. Figure 3.3.3 shows a DF image of orthogonal segments nucleating

fiom the Ge-platelet source. Note that the DF image dislocation lines display higher

resolution due to the isolation of the diffracted beam conditions achieved with the

objective aperture.

Figure 3 -3 -2: TEM BF image of CVD-9 after RTA for 5 s at 1 OOO°C

Figure 3.3.3: TEM DF image of CVD-9 after RTA for 5 s at 1000°C

3.4 Ion Imptantation

bLuch work has been published on ion implantation of heterostructures to snidy the

intermixing of quantum weki (Labrie er al., 1996 and Charbonneau et al., 1995) and the

formation of secondary and extended defect structures (Goldberg et al., 1995). Jaraiz er

ul. ( 1995) produced atomistic caiculations of ion implantation in Si. Their study of point

defsct and transient enhanced diffusion (TED) of dopants in Si is consistent with the now

îàmous -'+I" mode[, which States that the Si interstitial excess is assurned to equal the

implanted dose. It was aiso observed that d e r subsequent anneaiing, the fiee vacancies,

with their higher diffusivities, diffuse and annihilate most of the interstitials produced

during ion implantation. The dominant process being recombination. Some of the

vacancies reach the surface like a perfect sink and annihilate there, Ieaving behind an

excess of interstitials in the bulk materiai (Jaraiz et al., 1995).

In this study, the area of concem has been focused on the direct primary eff- of ion

implantation in the superiattice. As stated earlier, the Ge-rich platelets, which possess an

interstitial elastic saain field, are relieved by the injection of point defects via ion

implantation. These small Ge perturbations require the diBision of vacancies to produce

the incipient dislocation loop. The excess interstitiah, as stated above, make this pro-

more dificuit by "rnopping up" the vacancies left behind. The relationship between the

spatial distribution of the secondary defects and the original primary effeçts is unclear

(Goldberg et al., 1995).

Ion implantation range profiles are very easily calcdated using wmputer simulations

(Le. TRLM and binary collision approximation @CA)) due to the relative simplicity of

the collision processes. However, ion implantation damage after subsequent annealhg is

far more complex. Molecuiar Dynamics (MD) provides detailed predictions of the

damage but only for the first nanoseconds because of its heavy compuiational burden

(Jaraiz et al., 1995). It has been demonstrated in Section 3 2 that the primary effets of

ion implantation have assisted in the reduction of misfit dislocations by a vacancy

migration mechanism after RTA.

Secondary defect formation has also been observed in the cross-section TEM studies

(Figure 3.4.1). TRIM (Ziegler et al., 1985) calculations have been performed and the

depth profiles of the implanted ions have k e n produced for a 10-period Ge,&, (x =

0.09) (hsioe = 150 A, Hsi = 210 A) heterostnicture implanted with 540 keV Si ions,

sirnilar to the CVD-9 materia! of this shidy (see Appendk). The results were consistent

with those found in this TEM obsewation and aui be seen in the Appendix. The Si

implanted ions end-of-range lie between 6000 - 10000 A. This is the region where al1 the

secondary darnage can be observed. The top four layen of this CVD-9 cross-section

sample in figure 3.4.1 have been previously etched away in attempt to reveal the misfit

dislocations at the interface for Nomarski study.

After an 8j0°C RTA for 30 seconds, dislocation loops, dislocation dipoles and <3 1 1>

md dekcts were the structures making up the secondary damage. consistent with the

observations of Goldberg et al. (1 993). Figure 3.1.2 shows a plan-view image of the

same sample. In plan-virw, the depth of the implantation darnage c m not be determined

because from a top view, dl structures seem to be lying on the sarne plane. However, the

long orthogonal misfit lines running dong <110> directions can cleariy be distinguished

Figure 3 - 4 1 : TEM BF image of implantation profile in cross-section (z = [O 1 11) of CVD-9 RTA for 30 s at 850°C

from the rod-like defects and dislocation loops formed from the implantation and RTA.

Therefore. what has been introduced here is a very feasible solution for decreasing the

drnsities of misfit dislocations. The thin film of Ge,Sii., is what we are concerned with

in designing new electronic devices. The Si substrate wafer is of no use for the ultimate

device fabrication and will be discarded. Thus, the secondary defect stnictures present in

the substrate will not affect the outcome of the devices.

Figure 3.42: TEM plan-view image of implanted CVD-9 RTA for 30 s at 8jO0C

3.5 Low-Temperature MBE Heterostructures

The use of low temperature MBE has resulted in the irnprovement of coevaporative

doping control and the production of highly metastable, clastically strained

heterostructures. However, there exists a limit for the growth temperanire whereby

planar 2-D growth stops and the breakdown of the (100)-oriented growth occurs. Low

temperature growth of heterostructures have been examined in the past by Jorke et ai.

( 1989). They studied the breakdown of the epitaxiai growth at low temperatures as a

hnction of Si deposition rate for both constant and variable temperature growths. The

transition From single crystalline to polycrysralline (twinned) and finally arnorphous

growth was observed. Perovic et al. (1 991) studied the microvoid formation in low-

temperature MBE structures. They discovered the presence of small spheaicd structures

which accompanied the breakdown of the (100)sriented growth of thae superlattices via

the formation of { 1 1 1 aiented facets.

An interesting phenornenon has taken place in the low-temperature (- 400°C) MBE-

grown heterostnictures with the breakdown of (100)-oriented growth. TEM imaging has

been perfonned in cross-section to reveal the (100) surface of the superlattices. FE-SEM

secondas electron (SE) detector imaging has been performed to show the surface of

these structures and the orientation of the pits. Positron annihilation spectroswpy was

conducted as an additional study to observe the nahue of the S-parameter as a h c t i o n of

implant energy. An idea of the defect concentrations (i.e. voids) was observed fiom this

data.

Samples MBE-1688 and MBE-1690 were chosen for this study. These samples were

initially grown to closely resemble the structured geometry of CVD-9. However, with

the reduced growth temperatures, it was reveaied by a Shïmmel etch and Nomarski

interference microscopy that misfit disiocations were not being nucleated at the interface.

Therefore, other methods of examination had to be employed to detemiine what

mechankm was taking place. M e r reviewing some literature (Perovic et al., 1993),

Figure 3-51 wvas found. According to the growth parameters used to fabricate MBE-

1688 and MBE-1690, the formation of voids was eminentiy probable. Although Figure

3.5.1 shows the morphoiogicai instability phase diagram for Geo-xSio.7s/Si, an

approximate estimation for the behaviour of G e ~ & i ~ - ~ ~ / S i (Le. MBE-1688, -1690) was

made. For a total thickness of 360 nm for the entire superlattice, a growth temperature of

- 400°C, and a growth rate of 0.5 d s , void formation would be the result.

200 300 400 500 600 700 800 . 900 1 0 0

Temperature ('C)

Figure 3 S. 1 : Morphological instability phase diagram describing cntical thickness instabilitia as a function of growth temperature

FE-SEM SE-detector imaging was used to examine the surface of the samples. The

low-kV (i.e. 1 kV) experiment was used to obtain images with s m d interaction volumes.

Figure 3.5.2 shows a SE-detector image with the presence of pits on the surface of the

MBE- 1688 sample. It was found later, through TEM observation that these features were

actually surface cusps with { l l l}onented facets. Figure 3.5.3 shows the TEM cross-

section image. The faceted growth suiface can be cleady resolved here. The (100)

surface breaks down into a series of cusps and peaks, with each cusp having a pyramidal

structure bounded by ( 1 1 1 } facets. It has been demonstrated by Perovic et al. (199 1) that

if the Si epitaxial thickness is increased, cylindrical voids are le& in the wake of the

migrating surface at each cusp. These cylindrical channels become morphologically

unstable and break up to fom a senes of aiigned voids (i.e. microvoids). Due to the

small thicknesses of the Si spacer layers of MBE-1688 and MBE-1690 (Le. 2 1 nm of Si),

the spherical microvoids were never observed in these samples.

Figure 3-52: FE-SEM SE-detector image showing presence of pits on surface

Figure 3 3.3 : TEM cross-section image revealing cusps with ( 1 1 I }-oriented facets

The low growth temperatures encountered in the MBE faci!lty accounted for the

formation of the surface cusps with { I 1 1 } -0riented facets. If allowed to proceed, with

larger Si spacer thicknesses, microvoids would form in the wake of the migrating cusp.

At these Iow growth temperatUres, the thermal energy produceci is low enough to

decrease the surface migration energy of atoms and the proper bonding sites at the fke

surface are not reached. Vacancies agglomerate and form these cusp structures.

Figure 3.5.4 shows the positron annihilation spectroscopy data for MBE-1688 and

MBE- 1690 performed at the University of Western Ontario. The beam of monoenergetic

positrons is implanted into the sarnple where the positron loses its energy (thennalization)

and then diffiises through the soiid until annihilating either from this fieely diffushg state

or &om a trapped state (bulk defect or a surface state) as described in Section 1.5. Two

specimens were taken from each sample (i.e. an as-grown sample and an annealed

sample). It was observed here that the S-parameten of the as-grown samples both

indicate the presence of cusps due to the steep incline leading to the sharp peaks at

approxirnately 5 keV. Interestingly, the S-parameter is higher in the material grown at

higher temperature (Le. MBE- 1688). Alîhough the higher S-parameter indicates that the

vacancy concentration is higher, it may mean that the defect structure is slightly different

in the two samples. The fiaction annihilating at the surface (i.e. Fs see section 1.5) is the

dominant factor in the S-panuneter expression and yields the steep peaks in the S-

parameter of the as-grown materiai. The presence of the large cusps bounded by the

{ I 1 1 1-oriented facets resdted in a high Doppler shift giving a high S-panuneter. M e r

annealhg the samples at 950°C for 3 0 seconds, the S- parameters greati y decreased,

indicating that the vacancy structures diffirsed during the annealing pmcess leaving a

lower vacancy concentration behind as shown in positron annihilation data of Figure

3 -2.9.

a. 1 = 16a8 M8E as grown

a P 1690MBE A r r A A CIZA

I

Figure 3.5.4: Doppler-broadened annihilation line-shape parameter S as a function of positron energy

The S-parameter represented above displays a value of Sd = 0.56 for the defected

region at approximately 5 keV for MBE 1688. At higher energies (> 10 keV), al1

annihilations occur in the fieely diffusing state (Le. void-net, defect-fiee) and is

represented by an S-parameter value of Sr = 0.5 1. The ratio (Sr / Sd) is 0.91, which is

consistent with other reported values for silicon for monovacancies or divacancies,

impurity-vacanc y or inters titial complexes (Perovic et al., 1 99 1).

CHAPTER 4 - CONCLUSIONS AND RECOMMENDATIONS

The study of defect formation in low misfit Ge,Sii.JSi heterostruchtres grown by

UHV-CVD and MBE has been examineci and based on a quantitative @y, it has

been dernonstrateci that the onset of strain relaxation via misfit disiocations can be

controlled by point defect injection. Point defects, which are generated by ion

implantation, will diffuse during post-growth rapid thermal annealhg (RTA) and

subsequentiy relieve the stress concentrations surroundhg growth-induced Ge

perturbations.

These Ge-cich platelets, fouad only by straïn field contrasting in the TEM, were

previously assumed to be inherent features of MBE. This study has lead to their

discovery in CVD material as weli.

The activation energy for nucleation of misfit dislocations, which was determined by

large ara bulk measurements (Le. Nomarski Interference Microscopy), was found to

be Q. = 2.5 f 0.5 eV. This universally accepted value for nucleation of rnisfit

dislocations remained constant for as-grown and implanteci samples regardless of

growth method (i.e. MBE or CVD), indicating that misfit dislocation nucleation is the

rate-limiting step in the strain relaxation process.

The overall strain relaxation activation energy was found to be Q, = 4 2 f 0.5 eV for

the samples, again regardes of gtowth methud, which is consistent with combining

the nucleation and glide energy terms in series (i.e. Q, = Q. + Qv).

0 The low-temperature MBE growth experiments yielded a value of 0.5 f 0.05 eV,

which represented the energy associated with vacancy migration. With the large

supersaturation of vacancies encountered with the low-temperature MBE growths, the

nucleation of the dislocation loops would be solely controlled by vacancy migration if

vacancy aggregation to form srnail cluster nuclei had already occumd.

a The vacancy formation energy term was dropped. Accordingly, the overall strain

relaxation energy value also decreased to a value of 2.0 f 0.5 eV, indicating once

more that vacancy migration is the rate-lirniting step in the Iow-temperature MBE

growths.

r Without monitoring the growth rates in low-temperature MBE material, the growth

surface of these structures was no longer found to remain planar. Low-kV FE-SEM

SE-detector imaging revealed surface pits on the samples. A senes of cusps with

(1 l 1)-oriented facets were observed with cross-sectional TEM. If the Si spacer

layen were increased, these facets were assumeci to generate linear arrays of spherical

microvoids in the wake of their growth.

By combining rnacroscopic and microscopie analyses dong with positron

annihilation spectroscopy and ion implantation, the semi-empincal predictive model

for strain relaxation developed by Perovic and Houghton (1992, 1993) has been

thoroughly examined to describe the early stages of misfit stralli relief for Ge,SiiJSi

(100). The behaviour of misfit dislocations in the Stage I regime (misfit densities <

10' cm-2) was consistent wîth the predicted model. Even in the presence of a large

vacancy supersaturation, the results still remain consistent with the model.

Perhaps fiiture studies can be conducted to fully understand the entire relaxation

processes across al1 the stages (Le. Stage I - III). Multiplication and interaction

effects of dislocations would need to be investigated to explain these stages of seain

relaxation.

Finally, a novel approach to reducing misfit dislocation formation via ion

implantation has been investigated. This process can be incorporated into the

fabrication of these heterostnictures in the manufachuhg of electronic devices such

as heterojunction bipolar transistors (HBT) and light-emitting diodes (LED).

Effective Stress and Kinetic Mode1 (Houghton et ai., 1994)

The effective stress is determined by the imbalance between the resolved shear stress

acting on the t h d i n g dislocation slip system and the iine tension in the extending a/2

4 10> misfit dislocation segment The effective stress due to misfit strain, expenenced

by the threading a m of an a/2 4 IO>-type 60' dislocation of length h, assuming isotmpic

behaviour and equai elastic moduli, is given by:

where yl is the angle between the strained interface normal and the slip plane, B is the

angle between the dislocation iine and its Burgers vector, and R is the angle between the

Burgers vector and the direction in the interface, normal to the dislocation line. p is the

shear modulus, v Poisson's ratio, ~ ( h ) the in-plane strain at position h, H is the thickness

of the layer through wfuch the threading dislocation cut, Hi is the displacement of the ith

misfit dislocation fiom the fkee surface. The core parameter ,&? = 4 is assumed. The Si in

the second term represents the increase in misfit dislocation tension due to dislocation-

dislocation interaction for the case where Aa(h) AT is introduced to account for the extra

strain due to differentiai expansion on heating to the anneal temperature.

For a single GexSii,/Si mdtilayer structure with N periods of altemating GeSi and Si

of thicknesses h and H, and with strain relaxation at the first GqSii,/Si interface, the

effective stress r , ~ (in GPa) h m above may be expressed as:

0.55 ln 1 ON(h + X))

The misfit nucleation rate ( d 2 s - ' ) for any (100) oriented Ge&.JSi (x < 0.5)

heterostructure grown by MBE or CVD cm be summarized by the following semi-

empîrical relationship:

where No is the initial misfit dislocation source density present at t = O(i.e. substrate

threading dislocations, residual subshate-buffer layer precipitates etc.), k is the

Boltzmann constant equal to 8.62 x 10" eVK and T is the anneal temperature in Kelvin.

r , ~ is given in units of MPa respectively. Q, is the activation energy for nucleation, n and

B are material constants and ,K is the shear modulus.

Using the classic Orowan plasticity equation (dddt = N(t) V(I)b.d, where befl is the

effective Burgers vector magnitude of a 60' dislocation projected in the strained

interface. the overall straïn relaxation rate (s") is given by:

where al1 parameten are material constants is the shear modulus) or experirnentally

detennined for the given materials system. Qn + Qv represent the overoll strah relaxation

activation energy, Q,.

Enerw Balance A~proach (Perovic and Houghton, 1992)

To determine the nucleation of misfit dislocations, the first step in the process was to

consider the generation of the interfacial prismatic dislocation loops via the loss of

coherency at Ge-nch platelets. Weatherly (1 968), Ashby and Johnson (1 969) and Brown

and Woolhouse (1970) nrst developed the theory to explain coherency breakdown at

precipitates in metallic systems. (EJ, the elastic energy of a wherent platelet is baianced

against the self-energy of a dislocation loop (Ed) to get the foiiowing total energy Et = E,

- Ed:

where , p is the shear modulus, v is Poisson's ratio, e is the Naperian base, b is th

magnitude of the Burgers vector and &' is the unconseained lattice misfit strain between

the precipitate and rnatrix. R is the dislocation loop radius where the inclusion sûah

energy is a minimum. O is the correction factor for the approximate infuite body self-

energy. The dislocation core parameter, a I 1.

The next step is to consider whether these loops can lead to misfit dislocations at strained

layer interfaces. The total Helmholtz energy balance equation (2) was used. E = E, - Ed

f Es - TS, where Ed is the dislocation half-loop self-energy including the appropriate

correction factor O for a semi-infinite body and a = 1. Es represents the energy gained or

lost during the creation or removal of a surface step and TS is the entropy. The governing

equation for (100)-oriented interfaces in the diamond cubic slip system ((1 1 1 } 4 IO>)

can be written as:

The Ec term has been modifieci by replacing the average lattice misfit strain (G) by an

effective strain Gr which is a superposition of the strain field of the perturbation with that

of the strained layer. From equation (21, one cm calculate the critical radius for

nucleation (R') upon maximizing with respect to R. The activation energy for nucleation

(E') is obtained by substituthg R ' into (2).

Bulk Quantitative Raw Data

Misfit Dislocation Nudeation Rates (~rn-~sec-9

Trials Temperature (K-')

0.000786 O.OOO8 18 0.000853 0.00089 1 0.000932

MBE- 1 703 ( A s - ~ w I ~ ~

CVD-9 (As-grown) Triais

Temperature O<-') 0.000786 0.0008 18 0.000853 0.00089 1 0.000932

MBE- 1707 (As-grown) - Trials

Temperature (Ki) 0,000786 0.0008 18 0.000853 0.000891 0.000932

1

900 400 130 40 10

1

300 250 170 120 80

2

879 356 135 36 8

3

940 378 146 32 16

2

287 236 164 127 74

3

326 289 156 115 76

Misfit Dislocation Densities (cm-')

CVD-6 1 (Irnplanted) ,

t

Triais Temperature (K")

0.000786 0.0008 18 0.000853 O-O0089 1 0.000932

Trials Temperature (KL)

0.000786 0-0008 18 0.000853 0.00089 1 0.000932

Triais Temperature (RI)

0.000786 0.0008 18 0.000853 0.00089 1 0.000932

1

900 190 65 18 6

2

845 1 84 57 26 9

MBE- 1703 (AS-grown)

3

930 206 75 24 11

~tials Temperature (Ki)

0.000786 0.0008 18 0.000853 0.00089 1 0-000932

1

3 100 980 1 70 70 17

2

3087 954 156 73 14

3

3024 932 179 61 19

Low-T MBE MBE-1707 (As-gro~n)

Trials Temperature (TC')

0,000786 0.0008 18 0.000853 0.00089 1 0.000932

2

135 53 23 15 7

1

150 60 30 14 7

3

165 58 35 14 6

TRIM Simulation @y R Goldberg, UWO)

6000

Depth (A)

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