Electrical characterization of

impurity-defect complexes in

silicon

Naveen Goud Ganagona

Department of Physics

University of Oslo

A thesis submitted for the degree of

PhilosophiæDoctor (PhD)

2014 November

© Naveen Goud Ganagona, 2015 Series of dissertations submitted to the Faculty of Mathematics and Natural Sciences, University of Oslo No. 1643 ISSN 1501-7710 All rights reserved. No part of this publication may be reproduced or transmitted, in any form or by any means, without permission. Cover: Hanne Baadsgaard Utigard. Printed in Norway: AIT Oslo AS. Produced in co-operation with Akademika Publishing. The thesis is produced by Akademika Publishing merely in connection with the thesis defence. Kindly direct all inquiries regarding the thesis to the copyright holder or the unit which grants the doctorate.

Abstract

Thermal evolution of impurity-defect complexes in proton-irradiated mono-crystalline

silicon material has been investigated by deep level transient spectroscopy (DLTS).

Especially, the interaction between common impurities such as oxygen, carbon and

hydrogen, and intrinsic defects is addressed. Oxygen and carbon are introduced during

the materials growth while hydrogen occurs during device processing of Si wafers, like

fabrication of cells. The defect dynamics was investigated by post-irradiation annealing

of the samples, isochronally or isothermally, at increasingly higher temperatures up to

475◦C while measuring the concentration of the various electrically active defects.

Firstly, the kinetics of the transition from divacancy (V2) centers to divacancy-

oxygen (V2O) pairs was studied. A simultaneous transition of the donor and acceptor

states of V2 to those of V2O was established by applying optical DLTS together with

ordinary DLTS measurements. An experimental value for the diffusivity of V2 in the

neutral charge state has been deduced from the isothermal annealing data and the re-

sults seem to favor partial dissociation of V2 as the predominant migration mechanism

and may challenging an one-stage mechanism proposed in the literature. Further, firm

evidence for the identification of trivacancy (V3) and trivacancy-oxygen (V3O) deep lev-

els in the bandgap has been established, enabling data on the formation and annealing

kinetics of these complexes in both n- and p-type samples. In particular, the formation

kinetics of V3O has been studied in detail and experimental values of the migration

energy and diffusivity pre-factor of V3 have been determined.

Secondly, the annealing of the interstitial carbon-interstitial oxygen (CiOi) pairs

has been studied, and it is concluded that a dissociation is the prevailing mechanism.

The binding energy between Ci and Oi is estimated to be 1.7 eV. A correlated growth

of a defect level at 0.39 eV above the valence band edge (Ev) was observed with the

iii

disappearance of CiOi. Correlation with data from Photoluminescence spectroscopy

(PL) measurements on similar samples suggested that the defect may be a interstitial

carbon-interstitial oxygen dimer (CiO2i) complex.

Finally, an attempt to study vacancy-hydrogen related defects was made using

hydrogen implanted p-type samples and tentative evidence for a hydrogen-related deep

level center with an acceptor state located 0.45 eV below the conduction band edge

was formed.

iv

Dedicated to my parents ...

Acknowledgements

I would like to take the opportunity to thank everyone who supported me

to accomplish this work. First of all, I would like to express my sincere

gratitude to my main supervisor Prof. Edouard Monakhov for giving me

this opportunity and for guiding me through many research and scientific

problems. Thank you for your patience and effort in proof-reading the ar-

ticles and your feedback has been helped me to improve my writing skills.

Secondly, a large thank you to my co-supervisor Prof. Bengt Svensson for

his extraordinary supervision, interesting discussions and shaping our arti-

cles into high quality ones. Thirdly, thanks to my co-supervisor Dr. Lasse

Vines for his hands on guidance, countless day to day discussions and initial

corrections to articles.

I would like to thank all the colleagues of the MiNa-Lab for their co-

operation and discussions. Thanks to E-lab colleagues, Chi-kwong Tang,

Vincent Quemener and Helge Malmbekk, for their assistance, discussions on

DLTS and creating good working environment. Special thanks to Bahman

Raeissi and Augustinas for PL measurements. My office mates Bhoodoo

and Per Lindberg deserve a special mention for creating a friendly environ-

ment and outings. I would like to thank our lab engineers, Viktor Bobal

and Mikeal Sjodin, for their assistance especially Viktor for performing all

ion-implantations in this work.

I would also like to thank Vishnukanthan and Raja for their wonderful

company and making life comfortable in Oslo. I also must thank my friends

Satish, Vikas, Varsha and Amul Anganti for sharing personal stuff and joyful

trips.

Finally, I would like to thank all my family members who supported me

during this period, especially my mom (Vijaya) and brothers (GLN goud

and Praveen). This work would not have completed without the constant

support and companionship of my beloved wife (Navya) and I would never

ever enough thank her, love you forever!

iv

List of included papers

I. Formation of donor and acceptor states of the divacancy-oxygen

centre in p-type Cz-silicon

N. Ganagona, B. Raeissi, L. Vines, E.V. Monakhov and B.G. Svensson

Journal of Physics: Condens. Matter 24, 435801 (2012)

II. Transformation of divacancies to divacancy-oxygen pairs in p-type

Czochralski-silicon; mechanism of divacancy diffusion

N. Ganagona, L. Vines, E.V. Monakhov and B.G. Svensson

Journal of applied Physics 115, 034514 (2014)

III. Formation of single and double donor states of trivacancy-oxygen

complexes in p-type silicon

N. Ganagona, L. Vines, E.V. Monakhov and B.G. Svensson

Solid State Phenom. 205206, pp 213-217 (2013)

IV. Formation kinetics of trivacancy-oxygen pairs in Silicon

N. Ganagona, L. Vines, E.V. Monakhov and B.G. Svensson

Journal of applied Physics 116, 124510 (2014)

V. Identification of the carbon-dioxygen complex in silicon

N. Ganagona, L. Vines, E.V. Monakhov and B.G. Svensson

In manuscript (2014)

v

VI. Defects in p-type Cz-silicon irradiated at elevated temperatures

N. Ganagona, B. Raeissi, L. Vines, E.V. Monakhov and B.G. Svensson

Physica Status Solidi C 9, No. 1011, 20092012 (2012)

VII. PL and DLTS analysis of carbon-related centers in irradiated

ptype Cz-Si

B. Raeissi, N. Ganagona, A. Galeckas, E.V. Monakhov and B.G. Svensson

Solid State Phenomena Vols. 205-206 (2014) pp 224-227

vi

Contents

1 Introduction 3

1.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

2 Basic concepts in semiconductors 7

2.1 Properties of semiconductors . . . . . . . . . . . . . . . . . . . . . . . . 7

2.2 Point defects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8

2.3 Electrical properties of defects . . . . . . . . . . . . . . . . . . . . . . . 9

2.4 Defect Annealing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

2.5 P-N junction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

3 Experimental methods 15

3.1 Capacitance-voltage measurements . . . . . . . . . . . . . . . . . . . . . 15

3.2 Deep Level Transient Spectroscopy (DLTS) . . . . . . . . . . . . . . . . 16

3.2.1 DLTS spectra . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

3.2.2 Optical-DLTS and MCTS . . . . . . . . . . . . . . . . . . . . . . 23

3.3 Photoluminescence (PL) . . . . . . . . . . . . . . . . . . . . . . . . . . . 24

3.4 Ion Implantation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

3.5 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 28

4 Results and Discussion 31

4.1 The divacancy and divacancy-oxygen complexes . . . . . . . . . . . . . . 31

4.1.1 The divacancy (V2) . . . . . . . . . . . . . . . . . . . . . . . . . . 31

4.1.2 The divacancy-oxygen complex (V2O) . . . . . . . . . . . . . . . 34

4.1.3 Diffusion mechanism of V2 . . . . . . . . . . . . . . . . . . . . . . 37

4.2 The trivacancy and trivacancy-oxygen complexes . . . . . . . . . . . . . 39

4.3 Carbon-oxygen complexes . . . . . . . . . . . . . . . . . . . . . . . . . . 41

vii

CONTENTS

4.4 Hydrogen related complexes . . . . . . . . . . . . . . . . . . . . . . . . . 44

5 Conclusions and suggestions for future work 47

References 49

viii

CONTENTS

1

CONTENTS

2

Chapter 1

Introduction

1.1 Introduction

The discovery of the transistor by Bardeen, Brattain and Schockley in 1947, trig-

gered a revolution in semiconductor electronics. During the last 60 years there

has been an exponential expansion both in development and use of semiconduc-

tor devices. Semiconductors are omnipresent in applications such as computers,

satellites, solar cells and indeed in general electronics applications. Among the

different semiconductors, Silicon (Si) has been the dominant material for the

semiconductor industry due to its abundance, cheaper to produce silicon based

devices than other materials and superior processing properties. However, several

new materials replaced Si in some of the areas, due to its limitations in certain

applications. Especially, it is not well suited in (i) optical applications as it has

and indirect bandgap, (ii) high power and (iii) in high frequency applications.

The recent growth in the photo-voltaic (PV) industry together with the need

for specialized applications like radiation hard detectors, e.g. in the Large Hadron

Collider, has lead to renewed efforts in basic research in Si. Interestingly, while

electronic industry has been thriving reducing the size of devices, the PV in-

dustry has focused on reducing the material cost and increasing the efficiency.

Since the production of high quality Si is expensive, a low cost processes have

been developed to produce low quality Si often called as solar grade Si. This

creates a situation where a precise control of impurities and material quality is

decisive. The material properties are very sensitive to the atomic imperfections

in the crystal, so called point defects which are present in the material. The main

3

1. INTRODUCTION

sources for the point defects are: non-ideal material growth, high temperature

processing steps during device fabrication, doping techniques e.g. implantation,

spin-on-diffusion and laser doping, and/or irradiation. These defects are a serious

concern in most of the applications as they act as a recombination centers dete-

riorating the device performance, e.g reducing the carrier lifetime in photovoltaic

applications. Defects are much more pronounced when the material is exposed

to radiation. For example, radiation damage is a major concern in certain area

of semiconductor applications; including device operation in a space and for par-

ticle detectors in high energy experiments. On the other hand, irradiation is

a perfect tool in defect studies, since they can be created and the concentration

can be tuned in accordance with the limitations of the characterization technique.

Impurities present in the material can interact with intrinsic defects which re-

sult in higher order complexes, crucial for device operation. As such, high quality

material is generally a good starting point to investigate point defect reactions,

given that the defects can be introduced in a controllable and reproducible way.

Carbon (C) and oxygen (O) are the most common and dominant impurities in

Si, and are introduce during the growth. C and O are normally found on sub-

stitutional and interstitial site, respectively, and natively do not influence the

electronic properties of the material since they are electrically inactive. However,

Oi and Cs can form a wide variety of electrically active complexes with other

intrinsic defects such as vacancies and interstitials or together with other impu-

rities, e.g. thermal double donors. Hydrogen (H) is another common impurity

introduced during processing of Si wafers. H is known to passivate different types

of electrically active centers in Si, making it an important impurity in multi and

single crystalline silicon. However, theoretical and experimental studies show that

H can also form electrically active complexes from the interaction with vacancy

related defects.

In this thesis, thermal evolution of multivacancy-oxygen complexes and im-

purity related defects (O,C and H) has been investigated using, mainly, the Deep

level transient spectroscopy (DLTS) technique. In order to get more information

on some of the defects also Photoluminescence (PL) spectroscopy has been em-

ployed. The thesis is organized as follows: Chapter 2 introduces the basics for

understanding point defects and semiconductors relevant to the present study.

4

1.1 Introduction

Chapter 3 describes the experimental techniques used in the work. Chapter 4

summarizes the work and the main experimental results are disseminated in Pa-

pers I - VII. In particular, the results reported in Papers I and II focus on the

divacancy related defects. Papers III and IV point out the annealing results for

trivacancy and related complexes. The results reported in Papers V, VI and VII

focus on the investigation of annealing behavior of carbon-oxygen complexes and

formation of carbon-dioxygen complexes. Interaction of hydrogen with irradiation

induced defects will be discussed in section 4.4.

5

1. INTRODUCTION

6

Chapter 2

Basic concepts in semiconductors

This chapter introduces a brief overview of the fundamentals of semiconductor

physics to better understand the experimental methods and the results described

in subsequent chapters and attached papers.

2.1 Properties of semiconductors

Electrons in an isolated single atom can occupy discrete energy levels. In solids, a

large number atoms are brought close together, discrete energy levels merge into

so called energy bands which contains positioned packed energy levels.[1] The

occupancy of energy states by electrons is given by the Fermi-Dirac distribution

F (E) = 11+exp[(E−Ef )/kBT ]

where Ef is the Fermi level, kB Boltzmann constant and

T is the absolute temperature. [2] The energy bands are filled by charge carriers

or empty at sufficiently low temperatures. The uppermost energy band occupied

by electrons is called valence band (Ev) while the lowermost unoccupied band is

referred to as conduction band (Ec). Electrons in the conduction band are free to

move in response to an applied electric field and, thus, transport current. Solids

can be classified into three types depending on their ability to conduct an electric

current and they are: insulators, metals and semiconductors. In semiconductors,

the valence band and conduction band are separated by a band gap (Eg) which

is free of energy levels in an ideal materials. Electron can be thermally excited to

the conduction band and the empty state, ”hole”, left in the valence band can also

move, and lead to electrical conduction in semiconductors. In insulators, valence

and conduction bands are separated by a large band gap and no availability of

electrons even with an applied electric field. In metals, valence bands is overlap

7

2. BASIC CONCEPTS IN SEMICONDUCTORS

with conduction bands and free electrons are readily available. Conductivity of

semiconductors can be greatly enhanced by so-called doping in which specific

impurities are introduced in order to control the type of doping. By doping the

semiconductor with an impurity that donates an electron to the conduction band,

the material will have excess of electrons, and is called as n-type; while material

with excess of holes is referred to as p-type semiconductor.

2.2 Point defects

In an ideal crystal structure atoms are ordered periodically in a lattice. However,

defects always occur in the crystal. A point defect is an imperfection in a crystal,

basically characterized by one unoccupied lattice position or one interstitial atom,

molecule, or ion. Point defects exist in thermal equilibrium, but can in addition

be introduced during growth and processing. Point defects can be accumulated

under irradiation and/or implantation. The following section mainly focuses on

irradiation induced point defects.

Point defects by irradiation

When a semiconductor is subjected to irradiation with sufficient energy the elas-

tic collisions between the energetic particle and a native lattice atom can result in

a displacement of the atom out of its lattice site. As shown in Fig 2.1a, an extra

atom placed between the atoms in the lattice (interstitial (I)) and a vacant lattice

site atom (a vacancy (V)) together known as a Frenkel pair. If the transferred

energy between the particle and the atom is above a threshold value, it results

in a generation of a Frenkel pair. The threshold energy for defect production is

defined as the minimum energy to produce a Frenkel pair. For silicon, it is around

15 eV.[3]

The nature of the radiation damage mechanism depends on the energy and

type of particle. In Si, electrons, due to their small mass, need above 250keV in

order to transfer an energy equal to the threshold energy. The collision cascades

are more pronounced for heavy ions compared to light ions and electron irradia-

tion. An energy of several MeV is needed to create secondary collisions in electron

irradiation, while a few hundred keV is needed in the case of proton irradiation

which is mainly used in the present work. In secondary collision cascades, nearby

8

2.3 Electrical properties of defects

Figure 2.1: Frenkel pair formation a), annihilation b), and impurity interstitial c)

vacancies can pair together to form divacancy (V2)[4], trivacancy (V3)[5] etc.

In Si, only a few percent of the irradiation-induced vacancies and interstitials

survive as they annihilate themselves, as shown in Fig 2.1b.[6] The survived V’s

and I’s are highly mobile at room temperature and migrate through the lattice,

and can react with other impurities (X) that are present in the material, for

example X + I → Xi, as illustrated in Fig 2.1c. The resulting defects are crucial

since electrical properties of the material strongly depend on the defects. The

electrical properties of the defects will be discussed in the following section.

2.3 Electrical properties of defects

The electrical and optical properties of a semiconductor can be greatly influenced

by the presence of defects. In a perfect semiconductor, no energy levels exist

in the band gap. However, the presence of defects and dopants can introduce

unwanted energy levels. Defects can be categorized as either electrically active

or inactive. Electrically active defects form states in the forbidden band gap and

9

2. BASIC CONCEPTS IN SEMICONDUCTORS

can emit/trap electrons to/from the conduction band or valence band. If the

trap is negatively charged after it captures an electron then it is referred to as an

acceptor. If the trap is positively charged after emitting an electron then the state

is called as a donor state. Electrically active defects have different characteristics

depending on their level position in the band gap and they can be either shallow

or deep levels. Shallow states are located close to the valence band or conduction

band, which are generally used for doping, while deep states are usually further

away from the band edges and act as generation/recombination centers.

Emission and capture

Consider a point defect with an energy level (Et) in the band gap and uniformly

distributed with a concentration,Nt, throughout the semiconductor. Charge car-

riers, i.e, electrons and holes in the conduction and valence band respectively,

can interact with the defect state resulting in a filling and emptying of the defect

state. According to Shockley-Read-Hall theory, this is a statistical process.[7, 8]

The possible transitions are shown in Fig 2.2. The rate of capture of electrons

and holes is cnn and cpp (cn/cp are capture coefficient for electrons/holes), re-

spectively, while the rate of emission for electrons and holes from a defect state

is en and ep, respectively, where n and p are the concentration of electrons in the

conduction band and holes in the valence band, respectively. Taking all the re-

actions into account, the change in the fraction of the state occupied by carriers,

nt, as function of time, t, is given by

dnt

dt= (cnn+ ep)(Nt − nt)− (cpp+ en)nt (2.1)

The capture rate can be defined as

cn,p = σn,pνth(n,p) (2.2)

where σn,p is the capture cross section of an electron or a hole and νth(n,p) is

the thermal velocity. The thermal electron (hole) velocity is defined as νth(n,p) =√3kBTm∗

n,p, where m∗

n(p) is the effective electron (hole) mass and T is the absolute

temperature. Further, the emission rate for electrons ans holes from a level can

expressed by the following equation [9]:

en, p(T ) = νth(n,p)g0g1σn,pNc,ve

−ΔGkBT (2.3)

10

2.4 Defect Annealing

Figure 2.2: Emission and capture transitions mechanism of a defect level

where Nc,v is the effective density of states in the conduction/valence band

g1 is the degeneracy of an occupied state, g0 is the degeneracy factor of an non-

occupied state, ΔG= ΔH - TΔS is Gibbs free energy of the defect state, ΔH is

the enthalpy, and ΔS is the entropy. Depending on the values of the emission

and capture coefficients, different transitions will dominate, as shown in Fig 2.2

. It is also useful to define another parameter, apparent capture cross section

(σapp):

σapp = σneΔSkB (2.4)

By measuring the emission rate as a function of temperature T, we can extract

the activation enthalpy and capture cross section from the Arrhenius plot of

ln(en/T2) versus 1/T. ΔH is commonly interpreted as ΔG ignoring the entropy

factor in DLTS measurements (which is mainly used in this thesis). For deep

states in the band gap TΔS is usually assumed to be small, however, this is not

necessarily valid in all cases.

2.4 Defect Annealing

Defects are generally stable up to a certain temperature range and can undergo

several reactions. Information about the defect can be revealed by a careful ex-

amination of its annealing behavior, for example using different characterization

techniques (such as DLTS, Infrared (IR) absorption, PL etc.). Defect annealing

can often be divided into two processes: migration and dissociation. For mi-

11

2. BASIC CONCEPTS IN SEMICONDUCTORS

gration, the defects become mobile at a certain temperature and migrate until

captured at other impurities and/or defects that are present in the material. In

this process defects can form new complexes (such as V2+Oi→V2O [10]) with

higher thermal stability or they can annihilate, for example a vacancy can be

eliminated when it meets an interstitial. Dissociation of a defect occurs if the

annealing temperature is high enough to overcome the binding energy of the de-

fect, for example CiOi can dissociates into its components Ci and Oi, when the

temperature is above 350◦C ( CiOi→Ci+Oi).[11]

A first-order annealing kinetics process can be described by [12]:

dN

dt= −cN (2.5)

where c is the annealing rate constant and N is the defect concentration. A

characteristic of this behavior is that it involves only one process with a rate

proportional to the defect concentration. The solution to the above equation

is N(t)=N0e−ct, where N0 is the initial defect concentration. For a typically

thermally activated process the rate constant depends on temperature T and is

given by:

c(T ) = c0e−EakBT (2.6)

where Ea is the activation energy and c0 is the frequency factor.

From the frequency factor, an indication about the observed process can be

obtained, i.e, migration or dissociation. Dissociation typically gives a frequency

factor of c0 ≈ kBTh

while migration usually would result in c0 � kBTh

, where kBTh

is the frequency of the phonons in the lattice. (h is Planck constant) [13, 14]

Further, a diffusion limited reaction between two defects A and B with a

concentration of NA and NB, respectively, can be described by the following

differential equation

dNA

dt= −4πR(DA +DB)NANB (2.7)

where R is the capture radius for the reaction and DA and DB are the diffusion

coefficients for defects A and B, respectively. If only one of the two defects are

mobile at the temperature of reaction eg., DA �DB, then D = DA+DB ≈DA

and DA = D0Ae

−EakBT , where D0

A is the exponential pre-factor. In this case, the time

12

2.5 P-N junction

evaluation of NA and NB is generally not exponential and so-called second order

kinetics occur.

2.5 P-N junction

A general overview of pn-junctions will be given in this section which are build-

ing blocks for many electronic applications and in understanding of other semi-

conductor devices. A pn-junction is a two terminal device. We can obtain an

asymmetric pn junction when the concentration of the donors side(ND) is much

more pronounced than the acceptor concentration (NA) side or vice-versa. In

these junctions, charge neutrality implies that the depletion region extends into

the low doped region. In this thesis we have used asymmetric junctions in most

of the experiments. When two semiconductors, n-type and p-type, are joined

together, a large concentration gradient in charge carriers occurs at the junction,

which will lead to diffusion. Under thermal equilibrium conditions, highly mobile

electrons and holes from the n and p-type regions, respectively, diffuse across the

junction. Electrons leave behind positively charged donor atoms, and holes leave

behind negatively charged acceptor-atoms, thereby creating a space charge region

(W) at the junction (see Fig 2.5).

An electric field, built up around the p- and n-region, prevents the further

diffusion of charge carriers entering the junction, and W is often referred to as

a depletion region. A potential difference built-up around the junction is called

built-in potential (Vbi). The potential barrier can be modified by applying an

external bias (V) to the junction. The net current will be zero, the drift current

is exactly balanced by the diffusion current.

Assuming an abrupt junction and by applying the Poissons equation, one can

derive the following equation for the depletion width W [2]:

W =

[2ε0εr(Vbi − V )

q

(Na +Nd)

NaNd

]1/2(2.8)

where ε0 is the permittivity of free space, εr is the relative permittivity of the

semiconductor material, V is the applied voltage and is positive/negative for

forward/reverse bias, q is the elementary charge and Na/Nd is the doping con-

centration of acceptors/donors.

13

2. BASIC CONCEPTS IN SEMICONDUCTORS

Figure 2.3: a) Schematics of p- and n-type semiconductors with energy bands

and b) pn-junction with energy bands, depletion region, contact potential and

electric field -

The capacitance of the depletion region is given by

C =ε0εrA

W= ε0εrA

[q

2ε0εr(Vbi − V )

NaNd

(Na +Nd)

]1/2(2.9)

where A is the area of the junction cross section.

For an abrupt n+p junction (Nd �Na), the capacitance given in Eq. 2.9 can

be reduced to

C = ε0εrA

[q

2ε0εr(Vbi − V )Na

]1/2(2.10)

14

Chapter 3

Experimental methods

The different experimental techniques, Capacitance-voltage measurements (C-V),

Deep level transient spectroscopy (DLTS) and Photo-luminescence (PL) spec-

troscopy, which have been used in this thesis will be discussed in this chapter.

3.1 Capacitance-voltage measurements

Capacitance-voltage (C-V) measurements are a technique to characterize a p-n

junction or Schottky barrier junction. During the C-V measurements, a small

AC voltage typically in the range of 10-100 mV superimposed on a DC voltage

is applied to the structure at typical frequency range of 1 kHz - 1 MHz. The

resulting charge (Q) variation gives rise to a differential capacitance, which is a

measured quantity. The capacitance is defined by:

C =dQ

dV(3.1)

By rewriting the Eq. 2.10, we get

1

C2= − 2

qε0εrA2NA

V +2Vbi

qε0εrA2NA

. (3.2)

From the above equations, by plotting 1/C2 versus V, the concentration (in

this case NA) can be extracted from the slope and the extrapolation of 1/C2=0

gives Vbi. Moreover, from C-V measurements, information about the effective

charge carrier concentration as function of depth can be extracted by the following

15

3. EXPERIMENTAL METHODS

8 10 12 14 16 18 20 22 241.6

1.8

2

2.2

2.4

2.6

2.8

3 x 1013

Depth [μm]

Hol

e co

ncen

trat

ion

[cm−3

]

Figure 3.1: Charge carrier concentration profile in a p-type Si sample with a Schottky

barrier contact of Al. The profile reveals passivation of boron acceptors close to surface.

relation.

N(W ) = − 2

qε0εrA2

(ΔC−2

ΔVa

)−1

. (3.3)

Fig 3.1 shows an example from CV measurements, hole concentration versus

depth of a Schottky contact on p-type-Si. Hydrogen can unintentionally accumu-

late at the surface of the sample during the Schottky contact formation. Hydrogen

can passivate boron acceptors by forming neutral B-H complexes [15], and thus

a reduction in the net carrier concentration occurs near the surface, as shown in

Fig 3.1.

3.2 Deep Level Transient Spectroscopy (DLTS)

DLTS is a powerful technique that is widely used to study electrically active de-

fects in semiconductor materials. DLTS was introduced by D.V. Lang in 1974.[16]

The principle of DLTS is based on the measurement of capacitance transients due

to the emission of charge carriers trapped by a defect in the depletion region of

the material. From DLTS data, one can extract the activation enthalpy, apparent

capture cross section and concentration of the levels in the bandgap of materials.

16

3.2 Deep Level Transient Spectroscopy (DLTS)

A general DLTS principle of operation can be explained by using a Schottky

or pn-junction. An asymmetrical pn junction using a n-type substrate is shown

in Fig 3.2. Consider the junction when a deep donor level with an energy position

Et is present. When Et is below the Fermi level, the level is filled with electrons.

The junction initially starts with a reverse bias Vr, as shown in Fig 3.2. The

resulting depletion region has a width of Wr and all the traps above the Fermi

level are empty while the traps below the Fermi level are filled with electrons.

The width λ, which is given by√

2ε0εrq2Nd

(EF − Et), is the region between the reverse

bias width and crossing of the defect level and the Fermi level. A filling pulse Vf ,

forward bias or rather by decreasing the reverse bias, is then applied for a short

amount of time usually a few milliseconds in order to fill the traps. As shown in

Fig 3.2, the band bending decreases and traps are moved below the Fermi level

and being filled. During this time, the depletion region width reduces to Wf and

the corresponding capacitance is C0.

When the pulse is removed after a time tp, the reverse bias is returned to

its original value and opposite to filling occurs: the traps, which were below the

Fermi level, again move above the Fermi level, and will emit their trapped carri-

ers. In the first case immediately after removing the pulse, the depletion width

is more than Wr and the corresponding capacitance is Cr −ΔC, ΔC is propor-

tional to the amount of traps. All the filled traps start to emit their carriers and

with the time the capacitance approaches its initial reverse bias condition. This

gives a capacitance transient which is then recorded. The procedure is repeated

several times at each temperature interval in a temperature scan. The capaci-

tance transient in response to the voltage pulsing sequence is measured at each

temperature throughout. The temperature interval used in this thesis is 40-300K.

By assuming that Wr is greater than Wf and the amount of traps are small

ΔC � C, the capacitance transient can be expressed as

ΔC(t) =CrNt

2Nd

e−ent, (3.4)

if uniform depth distributions apply for Nt and Nd.

17

3. EXPERIMENTAL METHODS

1 3

2

2)

3)

1)

Vr

Vf

Cr

1

Figure 3.2: Illustration of DLTS principle of operation with three steps: 1) Reverse bias,

2) filling pulse and 3) charge carrier emission

18

3.2 Deep Level Transient Spectroscopy (DLTS)

Figure 3.3: DLTS spectrum resulting from the transient capacitance difference measured

from two xed times t1 and t2 at different temperatures. (box car weighting function)

3.2.1 DLTS spectra

A DLTS spectrum is a representation of the collection of capacitance transients

recorded after each filling pulse. Transients are measured consecutively as a

function of time at different temperatures, as shown in Fig 3.3. The DLTS signal

can be deduced by the following equation

S =1

ni

td+ti∑td

ΔC(t)ω(t) (3.5)

where ni is the number of measurement points of the capacitance transient

for the so-called i’th window, ti is the niτ td is the delay time and ω(t) is the

specific weighting function, several types of weighing functions are possible where

the so-called Box car is illustrated in Fig 3.3.

A deep level in the band gap will result in a peak in the DLTS spectrum. Since

the emission rate is temperature dependent, at certain T, the DLTS signal reaches

a maximum value. By varying the length of time to measure the capacitance

transient, the so-called time window, the peak shifts in the temperature scale

19

3. EXPERIMENTAL METHODS

140 160 180 200 220 2400

0.05

0.1

0.15

0.2

Temperature [K]

ΔC/C

rb

CiOi

5.1 5.2 5.3 5.4 5.5 5.6 5.7 5.8 5.9 6x 10−3

10−4

10−3

10−2

1/T [1/K]

e n/T2 (s

−1K−2

)

Et = 0.36 eVσna=2.2×10−15 cm2

Figure 3.4: (a) DLTS spectra for six different rate windows in the range of (20-640 ms)−1

and (b) Arrhenius plot for the defect level

20

3.2 Deep Level Transient Spectroscopy (DLTS)

and results in a set of DLTS spectra with different time windows. An illustration

of a DLTS spectrum with different time windows is shown in Fig 3.4a. A set of

values [en, T] are available for the peak maximum at the different windows. By

using the Arrhenius relation:

ln(en/T2) = ln(βσna)−ΔH/kBT (3.6)

where β is a constant factor, the slope of the line in Fig 3.4b is proportional to

the activation energy of the defect level and the extrapolated intercept at 1/T→0 gives the apparent capture cross section. The concentration of the deep level

can be extracted from Eq. 3.4.

Weighting functions

As mentioned earlier, several types of weighting functions are possible depending

on the required signal to noise ratio and the energy resolution. The main ones

used in this thesis were the lock-in and the GaverStehfest (GS4) functions. The

lock-in weighting function is a simple mathematical conversion and has a good

signal-to-noise ratio. The lock-in weighting function is given by:

ω(t) =

⎧⎨⎩−1 td ≤ t ≤ td + 2i−1τ

1 td + 2i−1τ ≤ t ≤ td + 2iτ

The GS4 weighting function was first suggested by Istratov in 1997 and derived

by Stehfest-Gaver algorithm.[17, 18] It gives better energy resolution with higher

capability of separating the peaks but has a poorer signal-to-noise ratio. The

GS4 weighting function is given as

ω(t) =

⎧⎪⎪⎪⎪⎪⎨⎪⎪⎪⎪⎪⎩

−1 td ≤ t ≤ td + 2i−2τ

25 td + 2i−2τ ≤ t ≤ 2i−1

−48 td + 2i−1τ ≤ t ≤ td + 3/2× 2i−1τ

24 td + 3/2× 2i−1τ ≤ t ≤ td + 2iτ

21

3. EXPERIMENTAL METHODS

50 100 150 200 2500

0.04

0.08

0.12

0.16

Temperature [K]

DLT

S si

gnal

arb

. uni

t

Lock−inGS4

Figure 3.5: DLTS spectra from a lock-in and GS4 weighting functions.

A difference between lock-in and GS4 is shown in Fig 3.5. The peaks are

broader for lock-in compared to GS4. Hence, GS4 can more easily resolve over-

lapping peaks, but at expense of an increased noise level.

DLTS is a very sensitive technique in terms of trap concentration as it can

detect defects with low concentrations less than 0.01% of doping concentration.

Hence, it can go down to 108cm−3 range in pure and high resistivity samples,

and it is a non-destructive technique. Another important feature of DLTS is

the possibility to obtain depth profiling of the defects. On the other hand, the

defect concentration should not be exceed 10-20% of the doping concentration for

the above mentioned assumptions to be valid. DLTS does not give structural or

chemical identity of the defects, and makes direct assignment of the defect levels

difficult. However, one can make comparison studies with other complementary

techniques such as infrared absorption (IR), electron paramagnetic resonance

(EPR), and photoluminescence (PL) to identify the defects.

22

3.2 Deep Level Transient Spectroscopy (DLTS)

3.2.2 Optical-DLTS and MCTS

DLTS is commonly used to study majority carrier defects and, thus, limited to

probe only parts of the band gap. Ideally, one would like to probe the entire

band gap which also includes information about the minority carrier defects. In

normal DLTS, only majority carriers are injected into the probing region where

as, in so called minority carrier transient spectroscopy (MCTS) both majority

and minority carriers are injected in to the probing region. This can be done by

using pn-junctions where minority carriers are injected electrically by applying

a forward bias pulse. A selective injection of minority carriers is done optically

during the filling sequence, sometimes called optical DLTS (ODLTS).[9]

In the technique based on forward biasing a pn-junction during the filling

pulse, both majority and minority carriers are simultaneously present in the in-

vestigated region, the probability of minority carrier occupation of levels is low for

the defects with higher majority carrier capture rates. In ODLTS, optical excita-

tion can be done either from back side or front side through the diode (Schottky or

pn junction). Back side excitation generate carriers in the bulk if the absorption

length is shorter than the sample thickness. Further, only minority carriers will

diffuse to the depletion region if the diffusion length is sufficiently large, while the

majority carriers diffusion is negligible due to the electrical field of the junction.

In front side excitation, both types of charge carriers will be generated close to

or within the junction and signicant contributions from majority carriers must

be taken into account. In the present work, back side optical excitation is used.

Fig 3.6 shows the comparison of DLTS, MCTS and ODLTS spectra from some

selected p-type samples. The DLTS spectrum reveals only majority carrier (in

this case holes) traps, H1 and H2. The MCTS spectrum is similar to the DLTS

one in addition to one minority carrier trap E1. On the other hand, ODLTS

analysis reveals two other minority carrier traps E2 and E3 which were not ob-

served by DLTS/MCTS in addition to E1. It shows the applicability of ODLTS

for identifying certain traps, and one can obtain a more firm identification of the

defect levels by combining both DLTS and ODLTS measurements. One of the

major advantages of ODLTS over forward biased-MCTS is that it can be applied

to Schottky diodes.

Fig 3.6 shows the comparison of DLTS, MCTS and ODLTS spectra from some

23

3. EXPERIMENTAL METHODS

100 150 200 250−0.01

−0.005

0

0.005

0.01

Temperature [K]

ΔC/C

rb

DLTSMCTSODLTSH1

E1

E2

E3

H2

Figure 3.6: Comparison between DLTS, MCTS and ODLTS spectra of proton irradiated

p-type CZ-Si

selected p-type samples. The DLTS spectrum reveals only majority carrier (in

this case holes) traps, H1 and H2. The MCTS spectrum is similar to the DLTS

one in addition to one minority carrier trap E1. On the other hand, ODLTS anal-

ysis reveals two other minority carrier traps E2 and E3 which were not observed

in DLTS/MCTS in addition to E1. It shows the applicability of ODLTS for iden-

tification certain traps. One can obtain a clear identification of the defects by

combining both DLTS and ODLTS measurements. One of the major advantages

of ODLTS over forward biased-MCTS is that it can be applied to Schottky diodes.

3.3 Photoluminescence (PL)

Photoluminescence (PL) spectroscopy is a contact-less and non-destructive method

to probe the electronic structure of materials and defects/impurities. The ba-

sic principle of PL is that the excitation light is absorbed by creating electron-

hole pairs, which then may recombine radiatively, giving away the excess energy

through the emission of light (photons). The sample exhibits luminescence via

photo-excitation, hence the term photoluminescence.

24

3.3 Photoluminescence (PL)

Electron-hole pairs are formed and subsequently recombine by one of the sev-

eral mechanisms as shown in Fig 3.7. The excitation causes electrons within the

material to move into permissible excited states. These electrons return to their

equilibrium states, the excess energy is released and may include the emission

of light referred to as a radiative process, or may not emit light through a non-

radiative process. The energy of emitted light (Photoluminescence) relates to

the difference in energy levels between the two electron states involved in the

transition between the excited state and equilibrium ground state.

Figure 3.7: mechanism of absorption and luminescence between valence and

conduction band (direct excitation) -

The advantages of PL are band gap determination, which is useful in par-

ticular when working with new compound semiconductors, and impurity levels,

recombination mechanisms, and material quality can be determined. Radiative

transitions in semiconductors involve localized defect levels. The photolumines-

cence energy associated with these levels can be used to identify specific defects,

and the amount of photoluminescence intensity can be used for the indication of

their concentration. Direct band gap materials give stronger PL than indirect

band gap materials such as Si.

As an example, from paper VII, Fig 3.8 shows the PL spectra of proton

irradiated p-type Si and after the heat treatment at 385◦C. The so called C-line

ascribed to the CiOi complex, dominates in the as-irradiated sample while the

intensity of the so-called P-line increases after heat-treatment at 385◦C concurrent

25

3. EXPERIMENTAL METHODS

Figure 3.8: PL spectra of p-type Si as-irradiated and after the heat treatment at 385◦Cfor 180 min

with a decrease in the intensity of the C-line.

3.4 Ion Implantation

Ion implantation has been widely used for introducing impurity profiles with pre-

cise control of the position and dose for many decades. Although the technique

did not enter mass production before the mid 1970s, Shockley issued a patent

already in 1954 with a detailed description of the relevant processes involved.[19]

During ion implantation, ionized atoms are accelerated by an electrical field and

bombard a target, for example a semiconductor wafer. When an energetic ion

enters a solid target, it starts to lose energy. The distance that the ion travels

in the semiconductor is called as ion range which can be projected into a cer-

tain direction, typically perpendicular to the surface (so-called projected range).

The projected range of the implanted ions can be manipulated by varying the

ion energy typically from about 1 keV up to several MeV. The energy loss in a

target material is a result of two mechanisms, electronic stopping and nuclear

stopping. Electronic stopping involves collisions between the incoming ion and

26

3.4 Ion Implantation

the target electrons. The interaction between the incident ion and a target atom

can be treated as a Coulomb scattering event. Nuclear stopping is due to the

elastic collision between the two constituents i.e, incoming ion and the target

atoms. Nuclear stopping causes atomic displacement (damage) in the target ma-

terial and post implantation annealing is normally required to restore the crystal

structure.

Implantation as Irradiation

The implantation damage, caused by the ions penetrating the material, can also

be advantageous for studies of defects in semiconductors, since the defect gener-

ation can be controlled by the energy, dose and type of implanted ions. When

light elements such as hydrogen are implanted, a measurable amount of point

defects can be formed. Here, we regard ion implantation as irradiation when

the implantation peak is far away from the probed region. In the present work,

irradiation done with 1.8 MeV protons gives rise to a projected range of ∼40 μm,

while the DLTS measurements probed depths of 1-2um. During the annealing

out-diffusion of H from the region around the projected range may be anticipated

but no indications of any H-related levels in the DLTS spectra were found.

In the studies of H-related defects, the implantation peak of hydrogen into Si

will be in the probing region. Implantation generates primary defects which in

turn will interact with each other as well as with hydrogen, generating multiple

secondary defect complexes. The damage generated from 1 MeV proton implan-

tation extends in principle all the way from the projected range (∼17 μm) to the

surface. However, in this case we have used an Al-foil in front of the sample to

reduce the penetration depth of the protons, more closely corresponding to the

DLTS probe depth.

High temperature proton irradiations were also used in some of the experi-

ments performed in this thesis. High temperature irradiations are useful in un-

derstanding the defect formation with higher thermal stability, in particular for

lifetime control of power devices.[20] This is done by, first, heating the sample to

the required temperature. Irradiation is then done at the elevated temperature

and the temperature fluctuations are not more than 10◦C. After the irradiation,

27

3. EXPERIMENTAL METHODS

Figure 3.9: Schematic structure of n+p-junction which is used in this thesis

the samples were cooled down to RT in order to minimize the annealing of defects.

3.5 Sample preparation

The samples used in the present work were mainly n+p structures and a short

description of the fabrication will be given here. A sample structure is shown

in Fig.3.9. CZ-Si wafers were cleaned with a standard RCA (RCA1-3) cleaning

procedure prior to oxidation. After this cleaning process, the wafers were dry

oxidized at 1100◦C for 3 hours to grow a 250 nm thick SiO2 layer. The oxide

formed on the back of the wafer is used as a protective layer during the diffusion,

to ensure that no dopants enter the backside. The protection of back side oxide

can be done by spin coating of photoresist.

Standard positive photolithography and wet etching using buffered oxide etch

(BOE) were then applied to open holes with a diameter of 200-2000 μm in selected

areas. When the photoresist is completely removed under the exposed regions, a

buffered oxide etch (BOE) is used to open the holes in the oxide. After the oxide

etch is complete, most of the photoresist can be removed by rinsing in acetone

for a few minutes.

28

3.5 Sample preparation

The n+ layer was formed by indiffusion of phosphorous (P) from gas phase in

a commercial tube furnace. One can also use ion implantation and subsequent

activation in a tube furnace or by rapid thermal processing. In order to improve

the electrical contact between the diode and a probe needle, an Ohmic contact is

formed by evaporation of Aluminum (Al) over the oxide opening, thereby capping

the hole with a metal as well. The thin oxide layer that has been formed over the

diode areas during the diffusion process is removed by HF cleaning. A BOE etch

is applied to remove the backside oxide. Al Ohmic contacts are then prepared by

thermal evaporation on the front side (n+ side), having a thickness of about 200

nm.

29

3. EXPERIMENTAL METHODS

30

Chapter 4

Results and Discussion

This chapter aims to give a brief background related to the studies that have

been done in the thesis. A summary of the obtained main results will be given

with suggestions for the future work.

The present chapter deals with three different point defect complexes. In the

first part, the most prominent intrinsic point defect stable at room temperature

(RT) is discussed, namely, the divacancy (V2), its derivative the divacancy-oxygen

(V2O) complex, and the diffusion mechanism of V2. Second part reveals the

recently identified trivacancy (V3) and trivacancy-oxygen (V3O) complexes, and

their detailed transformation kinetics. The third part is about the annealing of

carbon-oxygen complexes and formation of new defects with higher stability in

the range of 400-450◦C.

4.1 The divacancy and divacancy-oxygen complexes

4.1.1 The divacancy (V2)

Irradiation-induced vacancies and interstitials in silicon are highly mobile at room

temperature and can easily form higher order complexes or complexes with im-

purities present in the material, which are stable at RT. Among the higher-order

vacancy complexes, V2 is the most prominent intrinsic point defect stable at RT

and it is strongly enhanced in concentration after particle irradiation and ion im-

plantation. The V2 can form in two ways: direct displacement of two neighboring

Si atoms during the irradiation and secondly, by agglomeration of two migrating

single vacancies (V) which are produced in two different collisions. The diva-

31

4. RESULTS AND DISCUSSION

Figure 4.1: Visualization of the basic atomic structure of V2 with its six neighboring Si

atoms.

cancy structure can be regarded as two missing Si atoms next to each other and

is schematically shown in Fig 4.1. Initially, there are six broken Si-Si bonds

around the divacancy, however, four of these bonds (e.g., atoms 1-2 and 4-5, as

shown in Fig 4.1) form into two bent pair bonds.[21] The remaining two bonds

are broken, so called dangling bonds.

Extensive studies of V2 have been done by Watkins and Corbett in the 60’s

using primarily Electron Paramagnetic Resonance (EPR) measurements. It was

shown that V2 can appear in three different charge states: positive, neutral and

negative. Later, DLTS studies established a doubly negative charge state. Energy

levels of V2 observed by DLTS at about Ev+0.19 eV, Ec-0.43 eV, and Ec-0.23 eV

are attributed to positive (+/0), negative (-/0) and doubly negative (=/-) transi-

tions of V2, respectively.[21, 22, 23, 24, 25] In p-type silicon, only V2 (+/0) can be

detected by majority carrier DLTS, however, it has been reported that V2 (-/0)

can be observed in p-type samples by optical-DLTS (see section 3.2.3 for detailed

explanation).[22] In turn, in n-type material the acceptor states of V2, single and

32

4.1 The divacancy and divacancy-oxygen complexes

double negative, have been well studied by DLTS while V2 (+/0) is not observed

by MCTS and ODLTS due to its large electron capture cross section.[23] The

details of the energy levels of V2 are summarized in Table 4.1.

Charge state Energy

level (eV)

Comments

V2 (+/0) Ev+0.19 DLTS in p-type [22, 25]

V2 (-/0) Ec-0.43 DLTS in n-type[21, 23] and

ODLTS in p-type [22]

V2 (=/-) Ec-0.23 DLTS in n-type [21, 23]

Table 4.1: Details of V2 levels

The early EPR studies by Watkins and Corbett investigated the annealing be-

havior of V2 in both CZ- and FZ-Si materials irradiated with high doses of MeV

electrons (10−18cm−2). It was found that V2 is stable up to 200◦C and starts to

migrate beyond this temperature with an activation energy, Ea, of ≈ 1.3 eV.[21]

The study suggested that V2 is more stable in FZ-Si, as compared to CZ-Si, due to

lower impurity concentration. A dissociation process was suggested as the main

annealing mechanism of V2 in FZ-Si. The dissociation energy was estimated to

1.9 eV. In CZ-Si, V2’s anneal by migration and trapping at other impurities that

are present in the material. One of the most abundant impurities in CZ-Si is

interstitial oxygen (Oi) which is present at concentrations of 1017-1018cm−3. V2’s

are trapped by Oi, and the divacancy-oxygen interaction is normally the main

mechanism of the V2 annealing in irradiated CZ-Si, as other early studies have

suggested.[26, 27]

The annealing behavior of V2 in Si irradiated with low doses is also depen-

dent on other, less abundant, impurities. Previous DLTS studies showed that

hydrogen can influence the annealing of V2.[28, 29] However, the present section

deals with oxygen as the main impurity trap during the migration of V2 while the

section 4.4 presents the effect of Hydrogen on V2.

33

4. RESULTS AND DISCUSSION

4.1.2 The divacancy-oxygen complex (V2O)

The divacancy-oxygen (V2O) complex was first observed by Lee and Corbett in

an EPR study on heavily electron irradiated Cz-Si.[26] The defect was present at

pronounced concentrations after high electron doses with similar concentrations

to that of V2, and an EPR Si-A14 signal was assigned to V2O. The atomic struc-

ture is similar to V2 (Fig 4.1)except an extra oxygen atom breaking one of the

reconstructed bonds in one of the bent pairs of Si atoms.

Later, in positron annihilation spectroscopy studies on CZ-Si material, an in-

crease in the V2O concentration was observed after the annealing of V2. [30] It

was also found that V2O anneals out at ∼400◦C. Further, an absorption line at

833.5 cm−1 was ascribed to V2O.[31] Theoretical studies on the electrical states

of V2O complex have predicted one donor and two acceptor levels in the energy

band gap.[32, 33] The details of the energy levels of V2O are given in Table 4.2.

Charge state Energy

level (eV)

Comments

V2O (+/0) Ev+0.23 DLTS in p-type [25, 34]

V2O (-/0) Ec-0.46 DLTS in n-type [10, 35] and

ODLTS in p-type [34]

V2O (=/-) Ec-0.20 DLTS in n-type [10, 35] [10,

35]

Table 4.2: Details of the V2O levels

Later, in a DLTS study on irradiation induced defects in oxygen-enriched high

purity n-type FZ-Si, a gradual shift in the position of the V2(=/-) and V2(-/0)

peaks was observed during heat-treatment at 200-250◦C. This shift was attributedto concurrent formation of two new levels with positions at Ec-0.24 eV and Ec-

0.46 eV.[10, 35] These new levels have similar electronic structure to that of the

V2 levels and tentatively identified as the single and double acceptor states of

V2O. Further, it was found that the shift in the level positions occurs faster in

material with high oxygen content.

In DLTS studies by Trauwaert et al. on annealed electron irradiated p-type

34

4.1 The divacancy and divacancy-oxygen complexes

material using both CZ and FZ samples, a shift was observed in the position of

the V2(+/0) level to Ev+0.24eV.[25] The shift occurs faster in oxygen-rich CZ

samples and it was proposed that the new level is due to a transformation of V2

to V2O, and Ev+0.24eV is related to V2O(+/0). However, a clear identification

of the transformation of V2(+/0) to V2O(+/0) was lacking and a simultaneous

formation of the acceptor and donor states of V2O, using the same sample, have

not been found in the literature. Moreover, the kinetics of the transformation of

V2(+/0) was not studied quantitatively.

In paper I, a simultaneous transition of the V2 (+/0) and V2(-/0) states to

the corresponding states of V2O was demonstrated during heat treatment above

200-300◦C. By applying optical DLTS together with DLTS, a more complete un-

derstanding of the V2 and V2O related levels was obtained. During the isothermal

annealing in the range 200-300◦C, a gradual decrease in the V2(+/0) level is ob-

served and a new level at Ev+0.23 eV emerges, as illustrated in Fig 4.2. The

increase in the amplitude of the Ev+0.23 eV level exhibits a close to one-to-one

correlation with the loss of V2(+/0) and is consistent with the assignment to

V2O(+/0). Similar to V2(+/0), annealing of the V2(-/0) is peak observed in the

ODLTS spectra (Fig 4.2) and a new level occurs at Ec-0.46eV, assigned to V2O(-

/0).[10, 35, 36] Furthermore, the rate of transition from V2(-/0) to V2O(-/0) is

identical to the V2(+/0) to V2O(+/0) transition. These results provide firm evi-

dence for the previously assigned V2 and V2O levels. Here, one can notice that the

amplitudes of the V2(+/0) and V2(-/0) peaks are not the same. The amplitude

of the V2(-/0) (V2O(-/0)) signal is saturated for pulses >100 ms without reaching

the full strength of V2(+/0) (V2O(+/0)). This incomplete filling of V2(-/0) can be

explained by a balance between the capture rate of optically excited electrons and

the emission of the captured electrons. Under the present conditions the capture,

displaying a saturation time of >100 ms, is limited by the excitation power and

the proton-induced peak damage in the bulk of the samples (∼40 μm), which

suppress the amount of the excited electrons that can reach the probing region

from the illuminated back side. Accordingly, the capture rate becomes slow and

comparable with the emission rate at the temperatures where the V2 (-/0) peak

occurs.

35

4. RESULTS AND DISCUSSION

100 150 200 250−4

−2

0

2

4

6

8

x 10−3

Temperature [K]

ΔC/C

rb

Pre−annealed5min10min45min

BiCs

H(0.24)

CiOi

V2O(−/0)

V2(+/0) V2O(+/0)

VOV2(−/0)

Figure 4.2: DLTS (dashed lines) and ODLTS (solid lines) spectra of proton irradiated

samples after pre-annealing (200◦C/20min) and isothermal annealing at 300◦C for different

duration. The rate window used is (640 ms)−1, and the weighting functions are GS4 and

lock-in for DLTS and ODLTS, respectively.

36

4.1 The divacancy and divacancy-oxygen complexes

4.1.3 Diffusion mechanism of V2

V2 can diffuse large distances prior to dissociation. As mentioned earlier, V2

migrates at and above the temperatures of 200◦C and early EPR studies had

suggested that the activation energy for V2 diffusion is ≈ 1.3 eV [21]. In a recent

DLTS study by Mikelsen et al., the annealing kinetics of V2 and formation of V2O

have been revealed in electron irradiated n-type diffusion oxygenated float-zone

silicon (DOFZ-Si). [36] It was found that the formation of V2O occurs concurrent

with the annealing of V2 with an almost one-to-one proportionality within the

experimental accuracy, and the activation energy for both processes is equal to

1.3 eV. However, the transformation kinetics of V2 in p-type Si has previously not

been studied quantitatively. In this thesis, we obtained information about the V2

atomistic diffusion by maintaining the V2 to V2O transition in both p- and n-type

materials with a known concentration of Oi.

In Si, V2 cannot migrate by a single jump as there are no common nearest

neighbor sites to a pair of adjacent sites. Therefore, V2 can diffuse in two possible

ways: (i) a two-stage process, where the two adjacent vacancies have to separate

one lattice spacing and then re-unite, i.e., via partial dissociation, and (ii) an one-

stage process, where a jump of an atom takes place over a next nearest neighbor

atom.[37] EPR studies by Watkins and Corbett suggested that V2-diffusion occurs

via the two-stage process. Later, in theoretical results by Hwang and Goddard

[38], it was predicted that V2 migrates via an one-stage hopping of a Si atom

along the divacancy with a saddle point in the V-Si-V configuration. This dif-

fusion mechanism suggested based on the fact that no sizable barrier exists for

stabilizing the partially dissociated (V-Si-V) configuration. Calculations have also

suggested that the total energy difference between the V-V and V-Si-V states is

1.36 eV. Since the V-Si-V configuration is unstable, a full dissociation process

requires two Si atoms to move simultaneously between the two mono-vacancies.

In paper II, the annealing kinetics of V2 and its transition to V2O has been

studied in detail. Isothermal annealing in the temperature range 200-300◦C has

been performed using proton irradiated p-type CZ-Si samples. The annealing ki-

netics of V2 with corresponding formation of V2O is of first-order. The estimated

activation energy for this process is 1.3 eV. In the present experiments, the Fermi

level at the investigated temperatures (200-300◦C) is close to the middle of the

37

4. RESULTS AND DISCUSSION

1.6 1.7 1.8 1.9 2 .0 2.1 2.2

10−17

10−16

10−15

10−14

Diff

usiv

ity o

f V2 [

cm2 /s

]

1000/T [1/K]

From the loss of V2(+/0)

From the growth of V2O(+/0)

DV2 = 1.5x10−3 exp [−1.31(eV)/kT]

Figure 4.3: The temperature dependent of the diffusion coefficient for V2.

band gap i.e, below the V2(-/0) transition at Ec-0.43 eV and above the V2(+/0)

transition at Ev+0.19 eV. The extracted values are similar to those obtained in

Ref.[36] and reflects the neutral charge state of V2 in both cases. Applying the

theory of diffusion limited reactions, the diffusivity of V2 has been extracted and

is given by (1.5±0.7)×10−3 exp[-(1.31±0.03) eV/kT] cm2/s. The temperature

dependence of the V2 diffusivity is shown in Fig 4.3.

The above mentioned diffusion mechanisms of V2 have a direct impact on the

geometrical factor of D0. Assuming the frequency factor to be equal to the De-

bye frequency in Si [39] and exp(-ΔS/kb)≈1; the theoretical estimate for D0V2

is

4.8×10−3 cm2/s and 1.2×10−3 cm2/s for the one-stage and two-stage processes,

respectively. Hence, these first-order theoretical estimates of the pre-factor for

V2 migration seem to favor partial dissociation of V2 as the prevailing diffusion

mechanism, rather than a one-stage process, in the neutral charge state.

38

4.2 The trivacancy and trivacancy-oxygen complexes

4.2 The trivacancy and trivacancy-oxygen complexes

There has recently been an increased attention towards vacancy clusters (Vn).

Among the vacancy clusters, a definitive identification has been achieved only

for divacancy, while the other vacancy clusters have been studied to a lesser

extent.[40] Especially, the trivacancy (V3) and related complexes have recieved

a growing interest. V3’s can be introduced by irradiation of Si with high en-

ergy particles. In early EPR studies on neutron irradiated Si, the so-called A4

signal was assigned to V3 in a (110) planar configuration.[5, 41] In a combined

DLTS and DFT study by Markevich et al, it was found that V3 is a bi-stable de-

fect in two configurations, namely, fourfold coordinated (FFC) and (110) planar

configurations.[42, 43] V3 in the (110) planar structure gives rise to four charge

states in as-irradiated samples. The FFC structure is the lowest in energy for V3

with an acceptor state at Ec-0.075 eV. The energy levels for V3 in different charge

states are listed in Table 4.3.

Charge state Energy

level (eV)

Apparent capture

cross section

(cm2)

V3 (=/-) Ec-0.36 10−15

V3 (-/0) Ec-0.46 10−15

V3 (-/0) (FFC) Ec-0.075 10−15

V3 (+/0) Ev+0.19 10−15

V3 (2+/+) Ev+0.10 10−15

Table 4.3: Details of the V3 energy levels

In n-type samples, it has been observed that the transformation from single

and double acceptor states of V3 in the planar configuration to the stable FFC

structure occurs in the temperature range 50-120◦C.[42] It was also suggested

that the V3, like V2, is stable up to 200◦C where it starts to migrate and can be

trapped by other defects or impurities. Isothermal annealing of electron irradiated

n-type material in the temperature range 200-300◦C resulted in the disappear-

ance of the V3 levels and concurrent formation of two new levels at Ec-0.34 eV

and Ec-0.455 eV with an amplitude similar to V3. One of the most prominent and

abundant traps for V3 is interstitial oxygen atoms (Oi) and the new levels have

39

4. RESULTS AND DISCUSSION

Figure 4.4: Experimental and simulated DLTS spectra (GS4 weighting function) for the

Ev+0.11 eV (H(0.11)) and Ev+0.24 eV (H(0.24)) peaks for rate window of (80 ms)−1.

been assigned to doubly negative and singly negative states of the trivacancy-

oxygen (V3O) center. In p-type Si, a shift in the position of the V2 (+/0) peak is

observed, and it was discussed that the emerging level is a combination of V2O

(+/0) and V3O (+/0).[44] In the same study, a level at Ev+0.12 eV is observed

upon the disappearance of V3 (2+/+) and tentatively identified as V3O (2+/+).

An annealing study of electrically defects in proton irradiated p-type Si has

been performed in paper III. In parallel to the V2-V2O transition, formation of

two new levels, Ev+0.24 eV (H(0.24)) and Ev+0.11 eV (H(0.11)), is demonstrated

during isothermal annealing at 300◦C with one to one correlation in strength. A

clear energy seperation of these levels was made, as shown in Fig 4.4. The am-

plitude of these levels is about ∼30% relative to V2 in all the investigated sam-

ples. The results strongly support the identification of the levels Ev+0.24 eV and

Ev+0.11 eV as V3O (+/0) and V3O (2+/+), respectively. As these measurements

were done few weeks post irradiation, the levels related to V3 were not observed

in these samples. This can be explained by the transformation of V3 from the

40

4.3 Carbon-oxygen complexes

metastable (110) planar configuration to the stable FFC configuration.

In paper IV, further detailed studies on proton irradiated p- and n-type sam-

ples and the formation kinetics of donor states of V3O were undertaken. The

samples were irradiated at RT with 1.8 MeV protons to a dose of 5×1012 cm−2

and the stored at -20◦C until the initial DLTS measurements were carried out.

The divacancy levels at Ev+0.19 eV and Ec-0.44 eV and a new level with energy

position at Ev+0.10 eV were observed in the as-irradiated sample. After the pre-

annealing at 200◦C, a decrease in the Ev+0.19 eV level was observed, while the

peak at Ev+0.10 eV annealed out. The amplitude of the Ev+0.10 eV level was

equal to the decrease in Ev+0.19 eV peak. This observation is in agreement with

previously suggested single and double donor states of V3. The present results

confirm the assignment of the respective V3 levels (Table 4.3), and they disap-

peared after the thermal pre-treatment at 200◦C. The formation of V3O levels was

observed during the isothermal annealing in the range of 200-300 ◦C and it was

found that formation exhibited first order kinetics with an one to one correlation

with the V3 amplitude. The extracted value of the activation energy for V3O for-

mation is 1.50±0.04 eV with a pre-factor of (2.1±0.7)×1010 s−1. An experimental

value for the diffusivity pre-factor of V3 was deduced by applying the theory of

diffusion limited reactions and determined to be (8.5±3.5)×10−2 cm2/s.

4.3 Carbon-oxygen complexes

Carbon and oxygen are the most prominent impurities in Si. Understanding the

annealing behavior and identification of carbon and oxygen related defects has

been a significant part of the work presented in this thesis. Interstitial carbon-

interstitial oxygen (CiOi), also known as the K-center or the C(3) center, is an

important defect arising in irradiated Si. Interstitial carbon (Ci) is produced di-

rectly after the irradiation, and is electrically active with a donor level position

at Ev+0.27 eV and is stable up to 270K. At RT, Ci is therefore unstable and

migrating Ci’s are trapped by interstitial oxygen atoms, forming a CiOi com-

plex. CiOi has been identified by a several characterization techniques : FTIR

(absorption line at 862 cm−1 ) [11], PL (C−line positioned at 0.79 eV) [45] and

DLTS ( with an energy level about Ev + 0.36 eV ). CiOi is thermally stable up

to 300◦C. Interestingly, there is usually an increase in the concentration of CiOi

41

4. RESULTS AND DISCUSSION

before annealing out. In this regard, only a few reports have been found in the

literature, FTIR ([11, 46]) and PL [45] studies on electron irradiated Si material

with high doses. FTIR studies in Ref.[11] suggested that interstitials (In) are

released from other complexes and interact with substitutional carbon (Cs) in

samples containing high carbon concentration.

A detailed DLTS analysis of the annealing kinetics of CiOi is given in paper

V. The annealing of CiOi exhibits first-order kinetics except at initial annealing

times for temperatures in the range of 355-385◦C. The thermal activation energy,

Ea, of the process was found to be 2.5±0.2 eV, with a pre-factor of 7±5×1013

s−1. A pre-factor in this range is in good agreement with what is expected for

dissociation, and suggests strongly that this is the main annealing mechanism.

Based on these results and assuming the migration energy of 0.87 eV for Ci, the

binding energy was estimated to be ∼1.7 eV which is close to the value predicted

by spin density functional theory.[33, 47]

A significant increase in CiOi is observed during the initial annealing duration

and is studied in detail using another set of samples in the temperature range

325-355◦C. The details of samples are given in Table I in paper V. The con-

centration of CiOi remains constant after the pre-annealing at 300◦C for 45min.

A gradual increase is observed during the isothermal annealing. The increment

exhibits first-order exponential kinetics and the extracted values for activation

energy and pre-exponential factor were ∼1.94 eV and ∼8×1012 s−1. The present

observations corroborate those in Refs.[11] and [45] and modelling of the kinet-

ics using the theory for diffusion-limited reactions suggests that ICiOi complexes

form during the irradiation in significant concentrations.[48] These complexes dis-

sociate above ∼300◦C resulting in an initial increase in CiOi.

Formation of Ev+0.39 eV level

The formation of the so-called P-line positioned at 0.767 eV in PL spectra is

observed in carbon and oxygen rich material upon the annealing of CiOi. A

carbon-dioxygen (CiO2i) complex has been the suggested as the origin of the P-

line. However, the electrical activity and the annealing kinetics of this complex

are still unknown. Papers V, VI and VII in the thesis investigate the formation

42

4.3 Carbon-oxygen complexes

of the Ev+0.39 eV level in detail. In paper V and VI, a connection between

annealing of CiOi and the formation of Ev+0.39 eV level with detailed kinetics is

given. This is further supported by comparing DLTS and PL results in paper VII.

As mentioned earlier, CiOi dissociates and its constituents may react with

other impurities or defects. Boron-carbon and carbon-carbon related complexes

have lower thermal stability (below 300◦C). Among the other impurities, next

to Oi, the oxygen dimer (O2i) is the most abundant oxygen related complex in

oxygen rich Si. A reaction with the oxygen dimer was suggested as a mechanism

for the formation of a new level at Ev+0.39 eV. The formation rate was followed

by the annealing of CiOi and the formation kinetics is illustrated in Fig 2 of paper

V. From the kinetics, the growth of the new level follows an exponential increase

and then saturates. The O2i concentration is expected to be in the range of 1013-

1014 cm−3 in Cz-Si and the saturation of the Ev+0.39 eV growth is attributed

to the depletion of O2i’s. Density Functional Theory calculations by Ewels et al

[49] predict that CiO2i has a donor level with an energy level position at Ev+0.39

eV which strongly supports the present results. Moreover, in high-temperature

irradiated samples, the concentration of the new level was enhanced by 4-5 times

compared to that of RT irradiated samples (paper VI). It is known that irradiation

at elevated temperatures enhances the O2i concentration and this is a likely rea-

son for such a high concentration of the Ev+0.39 eV level after hot irradiation.[50]

In order to reach a more comprehensive understanding of the Ev+0.39 eV

level, PL studies were performed together with DLTS. In the experiments, the

same samples have been used for both PL and DLTS measurements, as discussed

in paper VII. The CiOi center (C-line), positioned at 0.789 eV dominates in the

PL spectra of as-irradiated samples. Formation of the Ev+0.39 eV level and the

so-called P-line positioned at 0.767 eV was observed in annealed samples with

disappearance of CiOi. The annealing behavior of the C- and P-lines in the PL

spectra were similar to the CiOi and Ev+0.39 eV levels in the DLTS spectra. The

P-line has previously been assigned as a carbon-dioxygen complex which provides

support for the suggestion of the Ev+0.39 eV level to originate from the CiO2i

complex.

43

4. RESULTS AND DISCUSSION

4.4 Hydrogen related complexes

Hydrogen is known to passivate different types of electrically active centers in

Si, and in particular interfaces and surfaces making it an important impurity in

multi and single crystalline silicon.[51] However, it has been shown that H can also

interact with the most abundant vacancy related defects, such as the vacancy-

oxygen pair (VO) [52] and V2.[29] As for the H related defects, Irmscher et al [53]

found that two levels with energy positions at Ev+0.27eV and Ec-0.32eV have

identical thermal stability and amplitude in H-implanted n-type Si samples. Fek-

lisova et al [54] have shown that electron irradiated p-type samples treated with

wet chemical etching also display the formation of the above-mentioned levels.

From the literature, two levels, Ev+0.27eV and Ec-0.32eV, have been attributed

to a donor level and an acceptor level of vacancy-oxygen-hydrogen (VOH), re-

spectively. [28, 53, 54, 55] Recently, also the simultaneous formation of these two

levels has been demonstrated using the same samples.[56]

There have been several reports on the influence of H on V2 and forma-

tion of new divacancy-hydrogen (V2H) related complexes. Theoretical studies

by Coutinho et al [57] have shown that V2H possess donor and acceptor states

with energy levels that are close to the respective levels of V2 and it was found

that V2H2 possess only one acceptor state. In the study by Bleka et al, [58] after

wet etching of electron irradiated n-type samples, it was found that the anneal-

ing of V2 correlates with the growth of a hole trap at Ev+0.23eV, tentatively

attributed to V2-H.

In the present work, we have made an attempt to identify V2-H related com-

plexes using DLTS and ODLTS. The samples used for these experiments were

p-type with a boron concentration of 1015 cm−3. To avoid non-uniform carrier

compensation, because H can passivate the boron acceptor and to get uniform dis-

tribution of H, multiple implantations of H were utilized. This was done through

an Al-foil (15μm) with implantation energies of 1080, 1120 and 1160 keV to a

dose of (1.5-5)×1010 cm−2, yielding a box-like H-profile extending up to ∼4m in

depth. As shown in Fig 4.5, all the defects are stable up to 100◦C. Thereafter,when H starts to be released from B-H complexes, annealing of certain defects

and formation of new defects are observed, as shown in Fig 4.5. In the presence

of H CiOi anneals out at lower temperatures compared to that in H lean samples

44

4.4 Hydrogen related complexes

Figure 4.5: DLTS (above) and Optical DLTS (below) spectra of H-implanted and an-

nealed Cz-sample. GS4 weighing function and a rate window of (640 ms)-1 was used.

45

4. RESULTS AND DISCUSSION

where it is stable up to ∼300◦C. Two dominant hydrogen related levels with en-

ergy positions at Ev+0.27 eV and Ec-0.32 eV arise after annealing at 125◦C in

the DLTS and ODLTS spectra, respectively, consistent with the assignment as

donor and acceptor states of the VOH complex. Two levels exhibit a simultane-

ous formation with one-to-one correlation. However, the amplitudes of these two

levels are not the same. Presumably due to the ’short’ filling pulse used in the

present experiments (100 ms) versus that in Ref 56 (500ms). A decrease in the

amplitude of V2(+/0) level is observed while a gradual shift in the V2(-/0) peak

position takes place in the temperature range of 125◦C. The emerging new peak,

labeled as E(0.45), has an energy level at Ec-0.45 eV (Fig 4.5) and is stable up to

around 175-200◦C. Further, theoretical studies predict that V2H has energy levels

close to the V2 levels. Involvement of H is expected in the formation of this new

defect, however, the identity is not known at this stage and further studies are

required.

46

Chapter 5

Conclusions and suggestions for

future work

The work in the thesis is devoted to the understanding of annealing and formation

kinetics of intrinsic and impurity related defects in silicon. Firstly, progress has

been made towards a thorough understanding of the kinetics of multivacancy-

oxygen (V2-V2O and V3-V3O) complexes and in particular, of the atomistic diffu-

sion mechanism of V2. Secondly, electronic level and kinetics of the CiO2i complex

have been elucidated where the identification of the CiO2i deep level is also sup-

ported by PL experiments.

The first part of the thesis is focused on the kinetics of V2 to V2O transfor-

mation and formation of V3O. As a result, quantitative values for the diffusivity

of V2 in the neutral charge state have been extracted over a wide range, and the

data allow us to discuss the V2 atomistic diffusion mechanism. The results appear

to favor partial dissociation of V2 as the prevailing mechanism for V2-migration.

Since only a few theoretical reports are available on the diffusion of V2, it would

be interesting and helpful with more theoretical support to understand the dif-

fusion of V2. Another interesting study would be to investigate the charge state

effect on the annealing of V2 and thereby on the diffusion of V2.

The results on proton irradiated samples in Papers III and IV have provided

firm evidence on the deep level identification of the V3 and V3O complexes. Fur-

ther, experimental values for the activation energy and pre-exponential factor of

the V3 diffusivity are obtained. However,to unambiguously identify the energy

47

5. CONCLUSIONS AND SUGGESTIONS FOR FUTURE WORK

levels ascribed to V3 and V3O, a simultaneous transformation study of the donor

and acceptor states from V3 to V3O using the same samples should be under-

taken. As the V3 center is relatively less studied, more theoretical data on the

diffusion mechanism of V3 would be helpful, and further experiments concerning

the charge state effect on V3 to V3O transition in highly doped material should

also be pursued.

The second part of the thesis is devoted to study the carbon-oxygen related

complexes, described in Papers V-VII, and more investigations regarding the

increase in the concentration of the CiOi center before annealing out are en-

couraged, together using complementary techniques mainly PL. Moreover, DLTS

studies including samples with well controlled and adjustable carbon and boron

concentrations should be also performed. A correlated growth of the Ev+0.39 eV

level with the annealing of CiOi is found and the formation of this level resembles

that of the P-line in PL spectra, corroborating the assignment as CiO2i. It would

be interesting to investigate the formation this level in samples with different

carbon and oxygen concentrations. FTIR studies on the CiO2i center should also

be pursued.

As discussed in the last part of thesis, there are many open questions still

remaining regarding the vacancy-hydrogen complexes. A preliminary study on

vacancy-hydrogen related complexes was made and we have observed a gradual

shift in the V2(-/0) peak to Ec-0.45 eV in the hydrogen implanted samples (section

4.4). Involvement of hydrogen in the Ec-0.45 eV defect is anticipated, however,

the exact identification is still unknown. A careful further DLTS study is needed

to conclude further the origin of this defect. More computational work should

also be devoted to the vacancy-hydrogen complexes.

48

References

[1] C.Kittel. Introduction to Solid State Physics. John Wiley & sons, Inc, 8th

edition,, 2001. 7

[2] B.G. Streetman and S. Banerjee. Solid State Electronic Devices. Pren-

tice Hall, 5th edition, 2000. 7, 13

[3] J. J. Loferski and P. Rappaport. Radiation Damage in Ge and Si

Detected by Carrier Lifetime Changes: Damage Thresholds. Phys.

Rev., 111:432–439, Jul 1958. 8

[4] J. W. Corbett and G. D. Watkins. Silicon Divacancy and its Direct

Production by Electron Irradiation. Phys. Rev. Lett., 7:314–316, Oct

1961. 9

[5] Young-Hoon Lee and James W. Corbett. EPR study of defects

in neutron-irradiated silicon: Quenched-in alignment under 110-

uniaxial stress. Phys. Rev. B, 9:4351–4361, May 1974. 9, 39

[6] B. G. Svensson, C. Jagadish, A. Hallen, and J. Lalita. Genera-

tion of vacancy-type point defects in single collision cascades dur-

ing swift-ion bombardment of silicon. Phys. Rev. B, 55:10498–10507,

Apr 1997. 9

[7] W. Shockley and W. T. Read. Statistics of the Recombinations

of Holes and Electrons. Phys. Rev., 87:835–842, Sep 1952. 10

[8] R. N. Hall. Electron-Hole Recombination in Germanium. Phys.

Rev., 87:387–387, Jul 1952. 10

49

REFERENCES

[9] P. Blood and J.W. Orton. The Electrical Characterization of Semicon-

ductors: Majority Carriers and Electron States. Academic Press, 1992. 10,

23

[10] E. V. Monakhov, B. S. Avset, A. Hallen, and B. G. Svensson.

Formation of a double acceptor center during divacancy annealing

in low-doped high-purity oxygenated Si. Phys. Rev. B, 65:233207, Jun

2002. 12, 34, 35

[11] B. G. Svensson and J. L. Lindstrom, J.m. Annealing studies of the

862 cm1 infrared band in silicon. physica status solidi (a), 95(2):537–

542, 1986. 12, 41, 42

[12] T. R. Waite. Theoretical Treatment of the Kinetics of Diffusion-

Limited Reactions. Phys. Rev., 107:463–470, Jul 1957. 12

[13] James W. Corbett. Electron radiation damage in semiconductors and

metals, chapter 7 In Solid state physics advances in research and applica-

tions. Academic Press, 1966. 12

[14] Michael Moll. PhD thesis. Radiation Damage in Silicon Particle De-

tectors. PhD thesis, University of Hamburg, 1999. 12

[15] J. W. Corbett S. J. Pearton and M. Stavola. Hydrogen in crystalline

semiconductors. Springer-Verlag,, 1992. 16

[16] D. V. Lang. Fast capacitance transient appartus: Application to

ZnO and O centers in GaP p+n junctions. Journal of Applied Physics,

45(7):3014–3022, 1974. 16

[17] A. A. Istratov. New correlation procedure for the improvement of

resolution of deep level transient spectroscopy of semiconductors.

Journal of Applied Physics, 82(6):2965–2968, 1997. 21

[18] A A Istratov, O F Vyvenko, H Hieslmair, and E R Weber. Crit-

ical analysis of weighting functions for the deep level transient

spectroscopy of semiconductors. Measurement Science and Technology,

9(3):477, 1998. 21

[19] S. William. Forming semiconductive devices by ionic bombard-

ment, April 2 1957. US Patent 2,787,564. 26

50

REFERENCES

[20] D.C. Schmidt, B.G. Svensson, J.L. Lindstrom, S. Godey,

E. Ntsoenzok, J.F. Barbot, and C. Blanchard. 2 MeV electron

irradiation of silicon at elevated temperatures: Influence on plat-

inum diffusion and creation of electrically active defects. Journal of

Applied Physics, 85(7):3556–3560, Apr 1999. 27

[21] G. D. Watkins and J. W. Corbett. Phys. Rev., 138:A543–A555, Apr

1965. 32, 33, 37

[22] C. A. Londos. Annealing Studies of Defects Pertinent to Radiation

Damage in Si:B. physica status solidi (a), 102(2):639–644, 1987. 32, 33

[23] B. G. Svensson, B. Mohadjeri, A. Hallen, J. H. Svensson, and

J. W. Corbett. Divacancy acceptor levels in ion-irradiated silicon.

Phys. Rev. B, 43:2292–2298, Jan 1991. 32, 33

[24] M. Asghar, M. Zafar Iqbal, and N. Zafar. Study of alpha radia-

tion induced deep levels in p-type silicon. Journal of Applied Physics,

73(9):4240–4247, 1993. 32

[25] M.-A. Trauwaert, J. Vanhellemont, H. E. Maes, A.-M. Van

Bavel, G. Langouche, and P. Clauws. Low-temperature anneal

of the divacancy in p-type silicon: A transformation from V2] to

VOy complexes? Applied Physics Letters, 66(22):3056–3057, 1995. 32,

33, 34, 35

[26] Young-Hoon Lee and James W. Corbett. EPR studies of defects

in electron-irradiated silicon: A triplet state of vacancy-oxygen

complexes. Phys. Rev. B, 13:2653–2666, Mar 1976. 33, 34

[27] L. J. Cheng, J. C. Corelli, J. W. Corbett, and G. D. Watkins.

1.8-, 3.3-, and 3.9- Bands in Irradiated Silicon: Correlations with

the Divacancy. Phys. Rev., 152:761–774, Dec 1966. 33

[28] K Bonde Nielsen, L Dobaczewski, K Goscinski, R Bendesen, Ole

Andersen, and B Bech Nielsen. Deep levels of vacancy-hydrogen

centers in silicon studied by Laplace {DLTS}. Physica B: Condensed

Matter, 273274(0):167 – 170, 1999. 33, 44

51

REFERENCES

[29] E. V. Monakhov, A. Ulyashin, G. Alfieri, A. Yu. Kuznetsov,

B. S. Avset, and B. G. Svensson. Divacancy annealing in Si: In-

fluence of hydrogen. Phys. Rev. B, 69:153202, Apr 2004. 33, 44

[30] A. Kawasuso, M. Hasegawa, M. Suezawa, S. Yamaguchi, and

K. Sumino. An annealing study of defects induced by electron irra-

diation of Czochralski-grown Si using a positron lifetime technique.

Applied Surface Science, 85(0):280 – 286, 1995. 34

[31] J.L Lindstrom, L.I Murin, V.P Markevich, T Hallberg, and B.G

Svensson. Vibrational absorption from vacancy-oxygen-related

complexes (VO, V2O, VO2) in irradiated silicon. Physica B: Con-

densed Matter, 273-274(0):291 – 295, 1999. 34

[32] M. Pesola, J. von Boehm, T. Mattila, and R. M. Nieminen. Com-

putational study of interstitial oxygen and vacancy-oxygen com-

plexes in silicon. Phys. Rev. B, 60:11449–11463, Oct 1999. 34

[33] J. Coutinho, R. Jones, P. R. Briddon, S. Oberg, L. I. Murin,

V. P. Markevich, and J. L. Lindstrom. I6nterstitial carbon-oxygen

center and hydrogen related shallow thermal donors in Si. Phys. Rev.

B, 65:014109, Dec 2001. 34, 42

[34] N Ganagona, B Raeissi, L Vines, E V Monakhov, and B G Svens-

son. Formation of donor and acceptor states of the divacancy-

oxygen centre in p-type Cz-silicon. Journal of Physics: Condensed

Matter, 24(43):435801, 2012. 34

[35] G. Alfieri, E. V. Monakhov, B. S. Avset, and B. G. Svensson.

Evidence for identification of the divacancy-oxygen center in Si.

Phys. Rev. B, 68:233202, Dec 2003. 34, 35

[36] M. Mikelsen, E. V. Monakhov, G. Alfieri, B. S. Avset, and

B. G. Svensson. Kinetics of divacancy annealing and divacancy-

oxygen formation in oxygen-enriched high-purity silicon. Phys. Rev.

B, 72:195207, Nov 2005. 35, 37, 38

[37] A. Seeger and K. P. Chik. Diffusion Mechanisms and Point Defects

in Silicon and Germanium. physica status solidi (b), 29(2):455–542, 1968.

37

52

REFERENCES

[38] Gyeong S. Hwang and William A. Goddard. Diffusion and dis-

sociation of neutral divacancies in crystalline silicon. Phys. Rev. B,

65:233205, Jun 2002. 37

[39] R. A. Swalin. Atomic Diffusion in Semiconductors, edited by D.

Shaw. Plenum, New York, page pp. 65110, 1973. 38

[40] M. Ahmed, S.J. Watts, J. Matheson, and A. Holmes-Siedle. Deep-

level transient spectroscopy studies of silicon detectors after 24

GeV proton irradiation and 1 MeV neutron irradiation. Nuclear In-

struments and Methods in Physics Research A, 457:588–594, January 2001.

39

[41] J.W. Corbett, J.C. Bourgoin, L.J. Cheng, J.C. Corelli, Y.H.

Lee, P.M. Mooney, and C. Weigel. Radiation Effects in Semiconduc-

tors, (Inst. Phys. Conf. Ser. 31 Bristol and London), 1976. 39

[42] V. P. Markevich, A. R. Peaker, S. B. Lastovskii, L. I. Murin,

J. Coutinho, V. J. B. Torres, P. R. Briddon, L. L. Dobaczewski,

E.V. Monakhov, and B. G. Svensson. Trivacancy and trivacancy-

oxygen complexes in silicon: Experiments and ab initio modeling.

Phys. Rev. B, 80:235207, 2009. 39

[43] J. Coutinho, V. P. Markevich, A. R. Peaker, B. Hamilton, S. B.

Lastovskii, L. I. Murin, B. J. Svensson, M. J. Rayson, and P. R.

Briddon. Electronic and dynamical properties of the silicon triva-

cancy. Phys. Rev. B, 86:174101, Nov 2012. 39

[44] V. P. Markevich, A. R. Peaker, B. Hamilton, S. B. Lastovskii,

L. I. Murin, J. Coutinho, V. J. B. Torres, L. Dobaczewski, and

B. G. Svensson. Structure and electronic properties of trivacancy

and trivacancy-oxygen complexes in silicon. physica status solidi (a),

208(3):568–571, 2011. 40

[45] G Davies, A S Oates, R C Newman, R Woolley, E C Lightowlers,

M J Binns, and J G Wilkes. Carbon-related radiation damage

centres in Czochralski silicon. Journal of Physics C: Solid State Physics,

19(6):841, 1986. 41, 42

53

REFERENCES

[46] LI Murin, VP Markevich, JL Lindstrom, M Kleverman, J Her-

mansson, T Hallberg, and BG Svensson. Carbon-oxygen-related

complexes in irradiated and heat-treated silicon: IR absorption

studies. 82-84 of SOLID STATE PHENOMENA, pages 57–62, 2002. 42

[47] D. J. Backlund and S. K. Estreicher. C4 defect and its precursors

in Si: First-principles theory. Phys. Rev. B, 77:205205, May 2008. 42

[48] N. Ganagona, L. Vines, E. V. Monakhov, and B. G. Svens-

son. Identification of the carbon-dioxygen complex in silicon. In

manuscript, 2014. 42

[49] C. P. Ewels, R. Jones, S. Oberg, J. Miro, and P. Deak. Shallow

Thermal Donor Defects in Silicon. Phys. Rev. Lett., 77:865–868, Jul

1996. 43

[50] JL Lindstrom, T Hallberg, D Aberg, BG Svensson, LI Murin,

and VP Markevich. Formation of oxygen dimers in silicon during

electron-irradiation above 250 degrees C. 258-2 of MATERIALS

SCIENCE FORUM, pages 367–372. TRANSTEC PUBLICATIONS LTD,

1997. 43

[51] B.L. Sopori, X. Deng, J.P. Benner, A. Rohatgi, P. Sana, S.K.

Estreicher, Y.K. Park, and M.A. Roberson. Hydrogen in silicon:

A discussion of diffusion and passivation mechanisms. Solar Energy

Materials and Solar Cells, 4142(0):159 – 169, 1996. 44

[52] G. D. Watkins and J. W. Corbett. Defects in Irradiated Silicon.

I. Electron Spin Resonance of the Si-A Center. Phys. Rev., 121:1001–

1014, Feb 1961. 44

[53] K Irmscher, H Klose, and K Maass. Hydrogen-related deep levels

in proton-bombarded silicon. Journal of Physics C: Solid State Physics,

18(23):4591, 1985. 44

[54] O. Feklisova, N. Yarykin, E.B. Yakimov, and J. Weber. On the

nature of hydrogen-related centers in p-type irradiated silicon.

Physica B: Condensed Matter, 308310(0):210 – 212, 2001. International

Conference on Defects in Semiconductors. 44

54

REFERENCES

[55] P. Leveque, null A Hallen, B. G. Svensson, J. Wong-Leung,

C. Jagadish, and V. Privitera. Identification of hydrogen related

defects in proton implanted float-zone silicon. The European Physical

Journal - Applied Physics, 23:5–9, 7 2003. 44

[56] H Malmbekk, L Vines, E V Monakhov, and B G Svensson. Hy-

drogen Decoration of Vacancy Related Complexes in Hydrogen

Implanted Silicon. Solid State Phenom., 178-179:192, 2011. 44, 46

[57] J Coutinho, V J B Torres, R Jones, S berg, and P R Brid-

don. Electronic structure of divacancyhydrogen complexes in sili-

con. Journal of Physics: Condensed Matter, 15(39):S2809, 2003. 44

[58] J. H. Bleka, I. Pintilie, E. V. Monakhov, B. S. Avset, and B. G.

Svensson. Rapid annealing of the vacancy-oxygen center and the

divacancy center by diffusing hydrogen in silicon. Phys. Rev. B,

77:073206, Feb 2008. 44

55