DEFECT STUDIES IN METALS, ALLOYS, AND OXIDES BY POSITRON ANNIHILATION SPECTROSCOPY AND RELATED TECHNIQUES
Sahil Agarwal
A Dissertation
Submitted to the Graduate College of Bowling Green State University in partial fulfillment of the requirements for the degree of
DOCTOR OF PHILOSOPHY
August 2021
Committee:
Farida Selim, Advisor
Abby Braden Graduate Faculty Representative
Alexander Tarnovsky
Marco Nardone
© 2021
Sahil Agarwal
All Rights Reserved iii
ABSTRACT
Farida Selim Advisor
Positrons administer a unique non-destructive approach to probe materials with atomic- scale sensitivity and provide reliable information about the nature and size of defects. This study reflects on the powerful capabilities of positron annihilation spectroscopy (PAS) to characterize point defects at atomic-scale, which can be crucial in the development of relevant material for a wide range of applications like nuclear reactors, medical sciences, optoelectronic devices, nanotechnology, polymers etc. The work presented in this dissertation aims to gain a fundamental understanding of the defect structures in three unique material systems: Fe metal,
Fe-Cr alloy, and Ce:YAG oxide by applying the methodology and concepts of PAS. A wide range of other complementary characterization techniques has also been employed to enhance the understanding.
The depth-resolved PAS was used to identify vacancy clusters in ion irradiated Fe and measure their density as a function of depth. PAS measurements uncovered the structure of vacancy clusters and the change in their size and density with irradiation dose. Combining with
TEM measurements led to discovering a novel mechanism for the interaction of cascade damage with voids in ion-irradiated materials.
The effect of Cr alloying on the formation and evolution of atomic size clusters induced by ion irradiation in Fe-Cr alloys was also investigated using depth-resolved PAS measurements.
Combining with atomic probe tomography (APT), a possible explanation for the long-standing question about the well-known resistance to radiation in Fe-Cr alloys was addressed. It was attributed to the stabilization of vacancy clusters around Cr atoms that act as indirect sinks for radiation-induced defects. iv
The final part of this work focuses on studying the role of defects on the luminescence properties of an important photonic material, Ce:YAG. The work reports an interesting mechanism that modifies and completely reverses the photoluminescence (PL) temperature- dependent kinetics. Further, it is shown that PL temperature-dependent kinetics can be controlled by modifying microstructure and engineering defects.
v
To the greatest of sacrifices made by my beloved Mother and Father
To the new beginnings with my lovely wife Hayley vi
ACKNOWLEDGMENTS
I want to express my deepest gratitude and thanks to my advisor Dr. Farida Selim as without her support and guidance I could not have made this far. She has been a great role model for me throughout this journey. Through her strong presence and kindness, she has touched my life and helped groom me into the researcher I am today.
I also would like to extend my thanks and regards to Dr. Alexander Tarnovsky, Dr.
Marco Nardone and Dr. Abby Braden for agreeing to be in my committee and giving their valuable time in reviewing my manuscripts and for their insightful suggestions.
I am extremely grateful to my peers and lab mates, Dr. Pooneh Saadatkia, Dr. Petr Stepanov, Dr.
David Winarski, Md. Minhazul Islam, Armando Hernandez, Noalick Aboa and Micah Haseman.
They supported me at every step, helped me learn lab techniques and assisted me whenever I needed help. I am also thankful to all my collaborators and co-authors for being so supportive.
I am forever indebted to my wife, Hayley, who stood by me in these final years and helped me see this through. Her love and support guided me through challenging times and kept me motivated.
I want to dedicate this graduate journey to my parents, Rajni and Anil and to all the sacrifices they have made for me at every step of my life. They along with my sister, Surbhi, have been my greatest support system in every possible way.
vii
TABLE OF CONTENTS Page
CHAPTER I. HISTORY AND OVERVIEW ...... 1
Introduction ...... 1
Application Of PAS ...... 3
References ...... 6
CHAPTER II. POSITRON ANNIHILATION SPECTROSCOPY FOR ATOMIC DEFECT
STUDIES ...... 14
Introduction ...... 14
Positron Sources...... 15
Radioactive Decay ...... 15
Pair Production...... 16
Positron Interactions With Matter...... 17
Thermalization ...... 17
Diffusion ...... 18
Trapping In Defects ...... 18
Positron Wave Function, Momentum And Lifetime Distribution ...... 20
Doppler Broadening Of Annihilation Radiation Spectroscopy (DBS) ...... 21
Positron Annihilation Lifetime Spectroscopy (PALS) ...... 24
Trapping Model And Interpretation Of PAS data...... 28
Variable Energy Positron Annihilation Spectroscopy (VEPAS) ...... 29
Positron Implantation ...... 30
Positron Moderation And Slow-Positron Beams ...... 31
Complementary Techniques To PAS Employed In This Work...... 32 viii
X-Ray Diffraction (XRD) ...... 33
Scanning Electron Microscopy (SEM) ...... 34
Transmission Electron Microscopy (TEM) ...... 35
Atomic Probe Tomography (APT) ...... 36
Photoluminescence (PL) Spectroscopy ...... 36
Thermoluminescence (TL) Or Thermal Stimulated Luminescence (TSE)
Spectroscopy ...... 38
References ...... 41
CHAPTER III. DEFECT STUDIES IN ION-IRRADIATED IRON ...... 45
Introduction ...... 45
Experimental Methods ...... 46
Material Growth And Preparation ...... 46
Ion Irradiation ...... 46
Transmission Electron Microscopy (TEM) ...... 47
Variable Energy Positron Annihilation Spectroscopy (VEPAS) ...... 47
Results And Discussions ...... 48
Radiation Damage And Positron Stopping Profiles...... 48
DBS Measurements ...... 49
Depth-resolved PALS Measurements ...... 55
Defect Density Calculations ...... 59
TEM Measurements And Calculation Of Cavity And Void Density ...... 63
Conclusions ...... 67
References ...... 68 ix
CHAPTER IV. DEFECT STUDIES IN ION-IRRADIATED FE-CR ALLOY ...... 72
Introduction ...... 72
Experimental Methods ...... 74
Material Growth And Preparation ...... 74
Ion Irradiation ...... 74
Variable Energy Positron Annihilation Spectroscopy (VEPAS) ...... 74
Atomic Probe Tomography (APT) ...... 75
Results And Discussions ...... 76
Radiation Damage And Positron Stopping Profiles ...... 76
Depth-resolved Doppler Broadening Spectroscopy (DBS) Measurements ...... 77
Depth-resolved Positron Annihilation Lifetime Spectroscopy (PALS) ...... 80
Depth-resolved Defect Density Estimation ...... 83
Atomic Probe Tomography (APT) Analysis ...... 88
Conclusions ...... 91
References ...... 91
CHAPTER V. DEFECT STUDIES IN CE:YAG ...... 98
Introduction ...... 98
Experimental Methods ...... 101
Sample Preparation ...... 101
X-ray Diffraction (XRD) studies ...... 102
Scanning Electron Microscopy (SEM) ...... 102
Photoluminescence (PL) Spectroscopy...... 102 x
Thermal Stimulated Emission (TSE) Spectroscopy ...... 103
Positron Annihilation Spectroscopy (PAS) ...... 103
Results And Discussions ...... 101
Characterization of Ce:YAG NPs ...... 104
Investigation Of TD-PL Kinetics And Trap-Assisted Luminescence ...... 114
Conclusions ...... 122
References ...... 123
CHAPTER V. SUMMARY AND CONCLUSIONS ...... 129
COMPLETE LIST OF REFERENCES ...... 131
APPENDIX A. PUBLICATIONS ...... 156
APPENDIX B. CONFERENCE CONTRIBUTIONS ...... 158 xi
LIST OF FIGURES Figure Page
1.1 Histogram depicting research involving positron being published worldwide as a
function of year...... 3
2.1 The decay scheme of radioisotope 22Na ...... 16
2.2 Positron model potentials for negative, neutral, and positive charged vacancy ...... 20
2.3 The vector diagram of the momentum conservation in the positron annihilation
process...... 22
2.4 Doppler-broadened 511keV peak with the S and W parameters defined intervals ...... 24
2.5 Schematic outline of a coincidence DBS experimental setup ...... 25
2.6 Illustration of positron annihilation lifetime spectroscopy ...... 27
2.7 Makhovian distribution profile of positrons at implantation energies of 5-35 keV in
Fe...... 31
2.8 Illustration of the X-ray diffraction with Bragg's law ...... 34
2.9 Description of electron microscopy techniques ...... 35
2.10 Schematic setup of a PL system...... 38
2.11 Schematic demonstrating a simplified TSE process ...... 40
3.1 Damage and positron implantation profiles in the Fe thin films ...... 50
3.2 Doppler broadening measurements for the reference and irradiated samples featuring
defect parameter S as a function of implantation energy and depth ...... 55
3.3 Summary of depth-resolved PALS experiments...... 58
3.4 The properties of ion-beam-induced defects in the thin films from PAS ...... 62
3.5 TEM images of the pristine sample ...... 64
3.6 Post-irradiation defect quantification from TEM micrographs of the irradiated xii
sample ...... 65
3.7 The schematic shows the pre-existing voids in the reference films and the overlap
between them, and the cascade induced by irradiation ...... 67
4.1 Damage and positron stopping profiles for Fe and FeCr alloyed films ...... 77
4.2 DBS measurements showing depth-resolved S profiles and diffusion lengths ...... 80
4.3 Defect lifetimes and intensities as determined from depth-resolved PALS
measurements ...... 83
4.4 Estimation of trapping rates and defect densities in alloyed samples ...... 87
4.5 APT measurements for Fe-18 Cr films ...... 89
5.1 Unit cell of YAG crystal ...... 98
5.2 XRD patterns of Ce: YAG nanophosphors annealed in various temperatures from
1000 to 1500°C ...... 105
5.3 Grain size of Ce: YAG nanophosphors annealed at various temperatures from
1000 to 1500°C ...... 106
5.4 PL intensity of Ce: YAG nanophosphors as a function of wavelength ...... 107
5.5 PL emission intensity as a function of wavelength ...... 108
5.6 TD-PL kinetics of NPs ...... 110
5.7 TD-PL kinetics of SC...... 111
5.8 TD-PL kinetics of TC ...... 111
5.9 TSE emission in Ce:YAG TC ...... 112
5.10 Diagram illustrating the mechanism behind TL emission and its contribution to PL
spectrum ...... 114
5.11 TD-PL kinetics in various Ce:YAG microstructures ...... 115 xiii
5.12 Comparison of thermoluminescence glow curves of 3m and 10m avg grain
size TCs ...... 117
5.13 TSE emission in various Ce:YAG microstructures ...... 119
5.14 GIPS measurements on Single-crystal, 3m avg grain size TC and 10m avg grain
size TC ...... 121
xiv
LIST OF TABLES Table Page
2.1 Different radioactive sources used for positron production ...... 16
3.1 Table of 퐿 + values (average positron diffusion length) and the two lifetime components extracted for each sample ...... 54
5.1 Table showing corresponding values of average positron lifetimes. Larger lifetime value indicates the presence of large defect clusters ...... 121
1
CHAPTER I. HISTORY AND OVERVIEW
Introduction
The positron, the antiparticle to the electron, possess the same mass and spin as that of an electron but has the opposite charge. It carries one positive unit of the elementary charge.
Historically, the existence of positrons was first predicted by Dirac in 1928 as an interpretation for the negative energy solutions obtained from the relativistic invariant wave equation [1, 2].
Later in 1932, Anderson experimentally observed them in his famous cloud chamber experiment where he traced the trajectory of a particle from the passage of cosmic rays, which, when subjected to a magnetic field, showed a curvature identical to that expected for a particle with the mass-to-charge ratio of an electron but in the opposite direction [3]. The first annihilation studies of a positron with electrons in the matter were conducted in the 1940s, and positron physics emerged as a new field of research. The discovery of energy and momentum conversation during an ⅇ− − ⅇ+ annihilation event led to the development of traditional positron annihilation spectroscopy (PAS) methods in the next few decades and are still primarily used today.
When a positron interacts with its antiparticle, the electron mutual annihilation results in the emission of two 511 keV gammas. The very first PAS technique to originate was the Angular
Correlation of Annihilation Radiation (ACAR) in 1949, following the discovery that the angle between the two 511 keV annihilation gamma rays deviated from being oriented in the exactly opposite direction [4]. Subsequently, in 1951, the lifetime measurements of positrons in gases laid the foundation for the development of positron annihilation lifetime spectroscopy (PALS)
[5]. In the following years, the emergence of high-resolution Germanium detectors allowed
Doppler Broadening Spectroscopy (DBS) measurements to be performed. In the next couple of decades, a great deal of research elucidated the change in positron annihilation characteristics 2 when positrons are trapped at point defects and especially led to the emergence of PALS and
DBS as non-destructive methods for material characterization for probing atomic defects.
The first efforts in developing the positron beam techniques can be traced back to
Madanski and Rasetti in 1950 where they tried to moderate the positrons from a radioactive source to lower energies to obtain a monoenergetic source of positrons. However, they were unsuccessful [6]. Finally, in 1958, the first measurements on the emission of secondary electrons from solid surfaces bombarded with low-energy electrons and positrons were made [7].
Nonetheless, it was not until 1973, when the pioneering work by Canter et al. led to the discovery of smoked MgO moderators and provided a breakthrough in the field [8]. Thus, canter et al. paved the way for future advances in developing intense positron beams and ignited the need to understand positron-surface interactions.
At the beginning of the 21st century, Selim et al. in 2002 [9], took a giant leap forward beyond conventional PAS methods by inventing Gamma-induced positron spectroscopy (GIPS), where positrons are created directly inside the material via pair-production from high-energy γ- rays and thus completely eliminating the unwanted background and source contributions from positron decay curves. In the last few decades, a great deal of research has been carried out to develop slow positron beams in laboratories worldwide. However, only a very limited number of projects have been reportedly succeeded so far to achieve PALS capabilities [10, 11]. The construction of a pulsed positron beam facility at BGSU is underway with the goal to realize and develop the first depth-resolved PALS system ever in North America as a novel probe for studying the surface, subsurface, and interface properties of materials and thin films.
3
Applications Of PAS
Recently, the focus of the fundamental research is being shifted to atomic dimensions in almost every aspect of science, which ultimately presents the need for a reliable and accurate atomic-scale probe. Thus, the prominence of PAS has increased tremendously in the past two decades, as shown in the histogram below (Fig. 1.1), with around 7000 publications in 2020 featuring positron research [12]. Today, PAS has made its presence felt in every major area of science for defect investigation with PAS being effectively employed in many physical systems like metals [13, 14], alloys [15], semiconductors [16-18], polymers [19], ceramics [20], zeolites
[21], biological systems [22-24], organometallics [25], nanomaterials [26] etc.
Figure 1. 1: Histogram depicting research involving positron being published worldwide as a
function of year. The dataset has been taken from web of science public library.
Defects play a crucial role in determining fundamental materials properties, thus regulating their applications, which has been an enormous driving force in developing new 4 materials. The emergence of PAS has led to several pioneering works over the years [27-29].
Much active research has been conducted in materials science, especially in metals and metal alloys, for their immense applications in nuclear sciences, including fission & fusion reactors etc.
[13-15]. Point defects like vacancies, voids, dislocations are often present in their lattices intrinsically and can often be introduced during fabrication. When exposed to adverse environments like radiation, corrosion, or ion implantations, these defect structures change their characteristics and critically contribute to the material's response in extreme environments. Thus,
PAS is an extremely valuable probe for investigating metal-based systems and can provide a fundamental understanding of the defect transport properties and help advance the next generation of nuclear materials [13-15]. This work discusses the depth-resolved defect distribution and their evolution in ion-irradiated Fe and Fe-Cr alloys in subsequent chapters.
Many PAS studies have been solely dedicated to understanding the physics of more complex heterogeneous systems like semiconductors or composite materials in general [16-18,
29-31]. Special attention has been given to PAS studies of oxides as atomic-level point defects play a crucial role in determining the transport properties of many oxide-based systems [29-33].
Most common of all, vacancy-related defects often alter physical properties by introducing additional energy levels in the bandgap, which then impacts these materials' optical and electrical capabilities. Recent advancements in materials engineering led to the development of new fabrication methods. Popular oxide materials have now been fabricated with different microstructural properties, which can be categorized into unique classes: nanophosphor [34, 35], single crystal [36], bulk phosphor [37], thin-film [38], and recently transparent ceramics [35, 39].
Transparent ceramics are emerging as materials of great interest for applications such as solid-state lasers, optical window materials, and scintillators [40-44]. They are considered better 5 options than single crystals because of their ease of fabrication and lattice defects. Over the years, several synthetic approaches have been reported to fabricate highly transparent ceramics, such as chemical methods and sintering methods. The latter has been further evolved into separate specialized techniques like vacuum sintering, microwave sintering, hot pressure sintering, hot iso-pressure sintering and spark plasma sintering which varies in growth temperature and environment [45-49]. However, defect characteristics are not well understood in these materials and dedicated PAS studies are scarce in the literature. Chapter-5 of this report will discuss temperature-dependent photo-luminescence (TD-PL) kinetics of five different microstructures of one oxide material, cerium doped yttrium aluminum oxide garnet (Ce:YAG).
Ce:YAG is an important photonic material known for its efficient broadband yellow emission due to the 5d to 4f transition of Ce3+ and has immense applications in the fields of white light- emitting diodes, phosphor screens, scintillation etc. [50-52].
The structure of the rest of this dissertation is as follows: Chapter-2 describes the theory and experimental techniques associated with PAS methods. In addition, the complementary spectroscopic methods applied with PAS in this study are also briefly discussed. Chapter-3 discusses the first depth-resolved defect distribution studies in ion-irradiated Fe metal, and observation of a novel mechanism for cascade-void interaction is revealed. Chapter-4 builds upon the previous chapter with PAS studies being extended to Fe-Cr alloys, and the origin of radiation resistance in these alloys has been investigated. Chapter-5 is dedicated to the defect 6 investigation in Ce:YAG oxide. The role of defects in induced TD-PL kinetics and the origin of trap-assisted luminescence is discussed by combining luminescence techniques and PAS.
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phosphor glass by screen-printing technology and its application in LED
packaging. Optics letters, 2013, 38, 2240-2243.
52. Chewpraditkul, W., Swiderski, L., Moszynski, M., Szczesniak, T., Syntfeld-Kazuch, A.,
Wanarak, C., & Limsuwan, P. Scintillation properties of LuAG: Ce, YAG: Ce and
LYSO: Ce crystals for gamma-ray detection. IEEE Transactions on Nuclear
Science, 2009, 56, 3800-3805.
14
CHAPTER II. POSITRON ANNIHILATION SPECTROSCOPY FOR ATOMIC DEFECT
STUDIES
Introduction
Positron annihilation spectroscopy (PAS) is a well-known technique that utilizes positrons (antiparticle to electron) to probe the open volume atomic defects in materials, such as vacancies and vacancy agglomerates [1]. It is a non-destructive method and allows measurements up to the sub-ppm level of defect concentration with high sensitivity (detection of
1 vacancy in 107 atoms). The experimental techniques are well-supported by theory since the positron states in the lattice can be calculated from first-principle calculations (e.g., reference
[2]). The application of PAS techniques can be classified based on the positron source used. In traditional laboratory-based systems, the positrons are emitted by radioactive decay (commonly used: 22Na) with a broad energy distribution that results into deeper penetration of positrons into the sample bulk, typically with a mean depth ranging from 10 - 100 m. However, monoenergetic positrons generated from positron beams (typically with implantation energies in the range 0- 35keV) can be used to probe surface and near-surface regions of a sample. The fundamental principle behind defect sensitivity of PAS is that positrons become localized in an open volume defect site lacking positive charge and stay trapped there until eventually annihilating with a nearby electron, with the emission of two 511 keV gammas. These gammas carry defect-specific information and characteristics pertaining to individual defects type which can then be analyzed using two conventional methods, namely Doppler broadening spectroscopy
(DBS) to measure electron momentum distribution around the defect site and positron annihilation lifetime spectroscopy (PALS) to record the lifetime of a positron lived in the defect 15 site. In this chapter, the theoretical basis and foundational principles behind positron annihilation techniques will be discussed.
Positron Sources
Positrons can be created by two different mechanisms, pair-production or radioactive decay. Radionuclides are the most commonly used positron sources for PAS. Alternative method entails sufficiently high energy photon-nucleus reactions to create positron-electron pairs energies for which greater than 1.02 MeV are required. In addition, nuclides with a proton excess can also provide positrons as an excess proton will decay into a neutron by the emission of a positron and a neutrino, however, use of this method has been very limited.
Radioactive Decay
The most popular method to create positrons is the use of a radioactive source that creates positrons with its decays. While other methods require the use of an accelerator or a nuclear reactor; this method does not, thus make it versatile and much more desirable for small-scale laboratories. Additionally, it is far less expensive to operate and is not limited by competing for beam time or maintenance shutdowns.
Table 2.1 lists some of the common positron emitting radionuclides. For PAS experiments, the positron source should ideally have a high yield, a suitably long half-life and positron emission should be accompanied by the near simultaneous emission of a gamma photon
(specifically for PALS). Thus, 22Na is widely used as a positron source because of its long half- life (t1/2 = 2.6 years), high positron yield of about 91 % and positron emission is accompanied by a 1.275 MeV photon. The radioactive decay is shown schematically in Fig. 2.1. 16
Table 2.1: Different radioactive sources used for positron production.
Isotope Half-Life Positron Fraction
22Na 2.6 y 0.91
58Co 71 d 0.15
11C 20 m 0.99
44Ti 47 y 0.94
64Cu 12.8 h 0.19
68Ge 275 d 0.86
57Ni 36 h 0.19
90Nb 14.6 h 0.54
55Co 18.2 h 0.81
Figure 2.1: The decay scheme of radioisotope 22Na
Pair Production
Positrons can also be generated within the material as pair-production using high energy gamma particles. When a high-energy gamma ray interacts with a nucleus of an atom, it results 17 in emission of positron-electron pairs. However, the incoming 훾 must at least have the energy equivalent of the rest mass of the ⅇ− − ⅇ+ pairs i.e., greater than or equal to 1.02 MeV.
Positron Interactions With Matter
The interaction of positrons with the electrons in a material ultimately leads to the annihilation process with one of the electrons. An energetic positron incident on a solid surface will either backscatter or will implant into the material. Upon entering the medium, rapidly lose all their kinetic energy and thermalize (usually in the order of a few picoseconds) due to non- elastic interactions. After thermalization, positrons freely diffuse through the material density while experiencing repulsion from the positively charged nuclei core until they become localized in an open volume defect site lacking positive charge and stay trapped there until eventually annihilating with a nearby electron, with the emission of gammas. In this process, the rest mass energies of the annihilated ⅇ− − ⅇ+ pair is converted to two 511 keV 훾-rays. These 훾-quanta carry information about the state of the positron-electron pair before annihilation. Thus, by studying the emitted 511 keV 훾-quanta it is possible to obtain useful information about the local electron states which is characteristic for positrons that trap into unique crystal defects making the measurement of the annihilation properties information a valuable method for defect studies in materials. The positron typically diffuses up to a few hundred nanometers prior to annihilation. The value depends on the electron density of the material and on the possible presence, and density, of low electron density positron trapping sites [4-6].
Thermalization
Positrons when injected into a material via a radioactive source, usually have kinetic energy of approximately 200 keV [7]. The kinetic energy is quickly lost to the host lattice through ionizations and core electron excitations, which mainly depend on the positron energy 18 and the nature of host material [8]. In the eV range, energy loss occurs mainly via electron–hole excitations and phonon emission, until the positron reaches thermal energy [8]. For positrons in the MeV range, the thermalization time after impact on the surface is only a few ps [9, 10].
Diffusion
Upon thermalization, positrons behave as free charge carriers in a delocalized state and their progression is dominated by electrostatic interactions with host lattice. The diffusion of positrons can be described with the use of the diffusion annihilation equation [10]:
휕 푛(푟⃗, 푡) = 퐷 훻2푛(푟⃗, 푡) − 휆 푛(푟⃗, 푡) 휕푡 + 퐵 (2.1)
Where, 푛(푟⃗, 푡) is the positron density at position 푟 and time 푡, 퐷+ is the positron diffusion constant and 휆퐵 is the bulk annihilation rate. The diffusion coefficient 퐷+ can be calculated from the semiclassical random-walk theory and is given by [2]:
푘 푇 퐷 √ 퐵 휏 += 푚∗ 푟 (2.2)
Where, 휏푟 is the relaxation time for scattering event, 푘퐵 the Boltzmann constant, 푇 the
∗ temperature and 푚 the effective positron mass. The positron diffusion length 퐿+푏 is defined as:
퐿+푏= √퐷+휏푏 (2.3)
Where, 휏푏 is the bulk positron lifetime, a characteristic of a given material. This classical approach to positron diffusion is based on several assumptions, i.e., positron scattering must be isotropic and quasi-elastic and the relaxation time approximation in the Boltzmann equation is assumed to be valid [2, 11].
Trapping In Defects
Positron implanted in a material delocalizes into a free Bloch state, however, if an open- volume defect center is encountered, the positron becomes can become localized at the site. 19
Positrons get trapped in localized states at the defect sites and annihilate with electrons in the local environment. Vacancy defects are the dominant type of imperfections that localize positrons; the missing nuclear charge creates a deep negative potential trap. The normally marked changes in annihilation characteristics for positron trapping at open-volume defects, compared to those annihilating from delocalized Bloch states, is a central feature of PAS techniques, the annihilation parameters for trapped state positrons are characteristic of the defect type [6, 9].
Positron trapping to point defects is extremely sensitive to the local charge of the defect with respect to the lattice. The interaction between positrons and a vacancy defect can be represented using simple square well potentials to elucidate the trapping potential of the missing atomic core and superimposed long range Coulomb tails to account for the local charge of the defect as shown in Fig. 2.2 for vacancy with neutral, single negative and positive charge states
[12]. The potential observed at negatively charged vacancies will result in an attractive force for 20 positron trapping, whereas the positively charged vacancy presents a Coulomb barrier and markedly reduces the probability of positron trapping.
Figure 2.2: Positron model potentials for negative, neutral, and positive charged vacancy
Positron Wave Function, Momentum And Lifetime Distribution
The wave function of the positron in a regular solid material can be calculated from the
Schrödinger equation [2]:
−ℏ 훻2휓 (푟) + 푉(푟)훻휓 (푟) = 퐸 휓 (푟) 2푚 + + + + (2.4)
Where,
푉(푟) = 푉퐶표푢푙(푟) + 푉푐표푟푟(푟) (2.5) and where 푉퐶표푢푙(푟) is the electrostatic Coulomb potential and 푉푐표푟푟(푟) is the effect of the electron-positron correlation potential in the local density approximation.
Over decades of research, there has been a great deal of theoretical effort given to solving the positron state 휓+ from the Schrödinger equation. The two most important experimental observables with respect to PAS are the momentum distribution technique and the positron lifetime technique. In the momentum distribution techniques, the energy (DBS) and angular distribution (ACAR) of the 훾-rays are measured. These measurements carry the information on the state of the electron-positron annihilation pair prior to annihilation. The positron-electron 21 momentum distribution requires a knowledge of how the electron wave functions overlap with the positron wavefunction. The positron lifetime study (PALS) utilizes the fact that in a lower electron density site the positron will be long-lived. The annihilation of ⅇ− − ⅇ+ pair results in creation of two characteristic 511keV gammas. The resulting pair carry defect-specific information exclusive to individual defects type which can then be analyzed using two conventional methods, namely Doppler broadening spectroscopy (DBS) to measure electron momentum distribution around the defect site and positron annihilation lifetime spectroscopy
(PALS) to record the lifetime of a positron lived in the defect site. The techniques are described below in more details:
Doppler Broadening Of Annihilation Radiation Spectroscopy (DBS)
In principle, the DBS technique is based on the measurement of the doppler shift of the characteristic 511 KeV 훾-rays emitted after the annihilation of an electron-positron pair. This
Doppler shift arises due to the slight difference in the momentum distribution of electron- positron pairs now of their decay. Upon annihilation, the momentum of ⅇ− − ⅇ+ pair is transferred to the generated 511 KeV 훾-rays due to the conservation of momentum. This leads to a small angle deviation from collinearity between the two 511 KeV 훾-rays and also induces a doppler shift. The momentum conservation in the ⅇ− − ⅇ+ annihilation process is illustrated by the vector diagram in Fig. 2.3. 푃1 and 푃2 are the emitted photon's momentum respectively, 푃 is the momentum of the center of mass of the positron-electron pair, with its parallel component 푃∥ and perpendicular component 푃⊥.
In the laboratory coordinate system, however, the two rays emitted deviates from collinearity with an angle given by: 22
푃 휃 = ⊥ (2.6) 푚푒푐
where 휃 = 180o, is the angle between the two rays in the laboratory frame of reference.
The Doppler shift is given by:
훥푣 푣 = ∥ 푣 푐 (2.7)
2 where the longitudinal center-of-mass velocity 푣∥ = 푃∥/2푚𝑒𝑐 , and the Doppler shift at the
2 energy 푚𝑒𝑐 is:
푣 퐸 푐푃 훻퐸 = ± ∥ = ± ∥ 푐 2 (2.8)
Hence, the line shape of the annihilation radiation reflects the momentum distribution of electrons in matter.
Figure 2.3: The vector diagram of the momentum conservation in the positron annihilation
process.
Since, upon entering a medium, positrons rapidly thermalize and lose all their kinetic energy, therefore, the momentum distribution of the annihilation radiation is mostly constituted by the participating electrons as opposed to the thermalized positrons, whose momentum is negligible. After thermalization, the diffusing positrons get trapped in defects and annihilate with the valence electrons that have low momentum, and thus, the width of the energy spread of 훾- 23 rays around 511 keV line is smaller. In contrast, in defect-free lattices, positrons have a higher probability of annihilating with the high momentum core electrons, and thus the energy distribution of 훾-rays around the 511 keV line is more extensive.
In general, the energy distribution of 훾-rays, is quantified by two line-shape parameters namely, S and W. If the electron is moving toward a detector, the detected 훾-rays energy is increased through the Doppler shift. As a consequence of positron trapping at defects, the positron wave function primarily overlaps with the wave function of the valence electrons of the surrounding atoms, leading to less Doppler broadening in the annihilation peak. Thus, high- resolution energy measurements of the 511-keV photons can be used to determine the increase in the fraction of positrons annihilated by valence electrons as a result of defects. The increase can be evaluated from the line-shape parameter of the 511-keV peak (the so-called S parameter). The
S parameter is usually determined as the fraction of counts accumulated in the central region, around the 511keV annihilation line as depicted in Fig. 2.4. Similarly, the shift arising from high momenta core electrons contributes in high doppler shifts from the 511 keV peak. This can be 24 represented by the ‘W’ parameter which is defined as the fraction of counts accumulated in the wing region, around the 511keV annihilation line (Fig. 2.4).
;
Figure 2.4: Doppler-broadened 511keV peak with the S and W parameters defined intervals.
Figure 2.5 outlines a conventional coincidence-DBS setup featuring a 22Na positron source sandwiched between the two identical samples and two high-purity Ge detectors (hpGe) with high energy resolution to record the annihilated 훾-rays. The detector signals are collected to a computer with an MCA, after which it is digitized. The extra hpGe detector in the coincidence setup imposes collinearity of the detectors, which ensures time coincidence, i.e., that the two annihilation photons need to originate from the same annihilation event and drastically improves 25 the peak to background ratio to about 104. This is important to collect chemical information from the counts in wing-regions of the DBS spectrum.
Figure 2.5: Schematic outline of a coincidence DBS experimental setup.
Positron Annihilation Lifetime Spectroscopy
In conventional 22Na source-based positron annihilation lifetime spectroscopy, the positron lifetime is measured as the time difference between the emission of the 훾 quantum representing +emission in the source and the 511 keV annihilation 훾-rays. This measurement is possible with 22Na because the 1.27 MeV 훾-ray is emitted almost simultaneously with the birth of positron in the 22Na source. Moreover, the thermalization time of positron implanted in a solid is only a few picoseconds, which is negligible compared with the positron lifetime. 26
The average electron density available to the positron at the defects is typically lower than in bulk, and correspondingly, the annihilation rate of positrons at defects is lower.
Therefore, the lifetime of a trapped positron in the presence of defects is usually longer than that in a perfect lattice. Similarly, the lifetime of positrons is highly sensitive to the defect size
(different annihilation rates) and increases with the increase in the size of each defect as the average electron density becomes scarce.
Conventionally, in a typical bulk measurement setup (Fig. 2.6 (a)), the lifetime of a positron (PALS) is measured by registering the two 511 keV photons serving as start and stop signals, detected via two scintillation detectors (commonly used, BaF2) coupled with the two photomultiplier tubes. The two detector signals are analyzed with analog nuclear instrumentation electronics featuring constant fraction discriminators to choose photons of correct energy, a time- to-amplitude converter, letting through only pulses spaced close enough in time, and a multichannel analyzer (MCA) connected to a measurement computer. The result is a histogram of annihilation events as a function of time differences between the start and stops signals, also known as the positron lifetime spectrum (푁(푡)) and can be expressed in terms of individual positron lifetimes (휏푖) and corresponding intensities (퐼푖) as described by the equation below:
퐼푖 푡 푁(푡) = 훴푖 ⋅ ⅇ (− ) (2.9) 휏푖 휏푖
Figure 2.6 (b) shows a typical lifetime spectrum. In general, the experimentally obtained spectrum is modeled by the sum of exponentials convoluted with the timing resolution function of the measurement setup and usually is approximated by a Gaussian function with FWHM (full width at half-maximum) of less than 250 ps. Besides, a background contribution resulting from the source is usually present and thus needs to be accounted for. 27
Figure 2.6: Illustration of positron annihilation lifetime spectroscopy. a) Outline of a positron
annihilation lifetime spectroscopy experimental setup, b) A typical PALS spectrum and
decomposition into individual lifetimes.
28
Trapping Model And Interpretation Of PAS data
A quantitative analysis of the defect parameters obtained from PAS data can provide important information about trapping parameters like trapping rates, size, nature and concentrations of defects, chemical information about the defect sites etc. All this knowledge can be extracted from the PAS data by the description of the positron trapping model given by
Bertolaccini and Dupasquier in 1970 [13] and later generalized by Frank and Seeger in 1974
[14]. The underlying assumptions of this model are no notable positron trapping before the thermalization process, defects are homogeneously distributed in the sample with no interaction among each other, and the wavefunction of the positron is completely localized at the trapping site. The model describes positron trapping in simple rate equations, which is directly proportional to the defect densities and is easily derived from the parameters obtained from both
DBS and PALS spectroscopy.
A simple case of the trapping model is a lattice with only a single kind of open-volume defect present. In such a system, after thermalization, the freely diffusing positrons will annihilate with an annihilation rate of ‘휆퐵’ in the defect-free region of the sample. In addition, a fraction of positrons will also get trapped in the open-volume defect ‘D’ with trapping rate ‘퐾퐷’ pending their annihilation with rate ‘휆퐷’. Then, the corresponding rate equations for the annihilation of positrons present in bulk (푁퐵) and trapped at defect site ‘D’ (푁퐷) can be expressed as follows:
ⅆ푁 (푡) 퐵 = −(휆 + 퐾 )푁 (푡) ⅆ푡 퐵 퐷 퐵 (2.10)
ⅆ푁 (푡) 퐷 = −휆 푁 (푡) + 퐾 푁 (푡) ⅆ푡 퐷 퐷 퐷 퐷 (2.11)
The solution for the equations (2.10) and (2.11) results in decay spectrum of positrons (퐷(푡)) as described in eq. (4). 29
퐼 푡 퐼 푡 퐷(푡) = 1 ⋅ ⅇ (− ) + 2 ⋅ ⅇ (− ) (2.12) 휏1 휏1 휏2 휏2
Where, parameters 휏1 and 휏2 denote the lifetimes of positrons annihilated in bulk and defect-site
D respectively, with corresponding intensities 퐼1 and 퐼2 and are related to the trapping parameters as follows:
퐼1 = 1 − 퐼2 (2.13)
1 휏1 = 1 − (2.14) 휆퐵+퐾퐷
퐾퐷 퐼2 = 1 − (2.15) 휆퐵−휆퐷+퐾퐷
1 휏2 = (2.16) 휆퐷
Similarly, the trapping model described above can also be applied to systems with two different positron defects and more.
Variable Energy Positron Annihilation Spectroscopy (VEPAS)
Conventional PAS methods described above are applicable to characterize defects in a wide variety of samples but required them to be bulk as positron emitted from a radioactive source penetrates deeper in the order of the order of 0.1–1 mm in solids and is unpredictable.
However, moderated positrons through a positron beam of variable energy can provide depth- resolved resolution and offer the possibility to study surfaces and thin-films. The idea of variable energy positron annihilation spectroscopy (VEPAS) is to perform PAS on solid-state samples using a controlled monoenergetic slow positrons which can be accelerated to the desired implantation energy using an electromagnetic transport system and high-vacuum conditions to convey the particles to the target sample implantation available with a slow positron beam.
VEPAS is thus a powerful tool in defect-depth profiling in solid-state sample and allows the measurements on thin film samples for defect structure identification. 30
Positron Implantation
Positrons implanted to a solid surface may undergo a number of different processes. They may be backscattered or may penetrate the sample. The positron penetrating the solid rapidly thermalizes by losing its kinetic energy. The thermalized positron would then diffuse in the sample until annihilating with an electron. Mono-energetic positrons implanted into a sample fall on a depth profile that can generally be described by the Makhovian distribution [15].
푚푧푚−1 푧 푚 푃(푧, 퐸) = 푚 ⅇ [− ( ) ] (2.17) 푧0 푧0
The parameter 푧0 is a function of incident positron energy, given by:
푧̅ 푧0 = 1 (2.18) 훤[ +1] 푚 where 푧 is the mean stopping depth and 푚 is a material dependent parameter with empirical value 푚 approximately equals to 2. The relationship between the mean depth with the initial energy is assumed to be a power law:
퐴퐸푛 푧̅ = 휌 (2.20) where the constant 퐴 is found empirically to be 3.6 g/cm2 keV-1.6 and n = 1.6 are independent empirical parameters related to theoretical and experimental uncertainties, and 휌 is the material density. 31
Figure 2.7. Makhovian distribution profile of positrons at implantation energies of 5-35 keV in
Fe.
Positron Moderation And Slow-Positron beams
Moderation of fast positrons is accomplished via ‘moderator’ materials that has negative work-function for positrons. The fast positrons are passed through the moderator where they undergo thermalization and a small fraction of them re-emerges with a kinetic energy equal to the work-function of the moderator material. In past, several materials have been tested for the applications in positron moderation, of which thermally annealed W has become the most popular choice owing to its decent moderation efficiency and easy usability [16]. In a typical laboratory setting, a thin (~1 μm) film of Tungsten is placed on top of the radioactive nuclei, and the moderated mono-energetic positrons are generated. The moderation efficiency of W is 32 typically only 10−3 – 10−2 % for most moderator materials. The slow positrons upon moderation can be separated from the non-moderated positrons via use of a velocity selector which can placed just after the moderation stage.
The resulting DC beam from can then be modulated using electro-magnetic fields and the beam can be compressed into short bunches. A periodic acceleration and deceleration lead to its compression in time at a certain focal point. The arrival time of the bunch at the target is then provided by the bunching electronics. The idea of time bunching of slow positrons was first discussed by Mills in 1980 [17] but there are only very few operational pulsed slow positrons beams due to their complexity in construction and optimization. The basic components of a pulsed positron beam system consist of a reflection type chopper, a sub-harmonic pre-buncher and a double harmonic buncher. The function of the reflection type chopper is to generate a pulsed beam. The chopped beam is then compressed to 2 ns pulse width at the buncher by a sub- harmonic pre-buncher. Finally, the double harmonic buncher compresses the pulse width to about 150 ps at the sample.
Complementary Techniques To PAS Employed In This work
Over the last five decades, PAS has been one of the most reliable methods to probe open- volume vacancy defects with high precision and accuracy. The significant advantage of PAS over other methods is that it can probe point defects in a lattice beyond the resolution of conventional microscopies. However, the true strength of the PAS technique lies in its versatility to be used in a complementary fashion with other characterization tools. While PAS can be valuable in monitoring atomic-scale defects, combining it with other techniques that gauge physical parameters like structural, optical, and electrical properties can provide a fundamental understanding of the material characteristics and open new avenues for research and 33 development. Combining the atomic-scale resolution of PAS with the mesoscale resolution of electron microscopies can be extremely valuable for a complete picture of materials' microstructural properties. In this section, the principle behind some of the popular characterization techniques is briefly described below. The applications and their usage in this work will be subsequently discussed in the chapters to follow.
X-Ray Diffraction (XRD)
X-rays are electromagnetic waves with a wavelength of 10-9 – 10-10 m, which corresponds to the dimension of the atomic plane interspacing in the crystalline lattices. XRD is a cheap, non- destructive crystallographic method and thus a very popular technique to characterize materials
[18]. It can accurately provide structural information about material properties like lattice parameters, chemical composition, growth orientation, phase purity, defects information, and layer thicknesses of thin films, and much more [18-22].
The principle of XRD can be explained by Bragg's reflection, as depicted in Fig. 2.8. The ordered crystalline structures diffract the incident X-ray beam, and constructive interference occurs between the scattered X-rays when the path difference 2 ⅆ 푠푖푛 휃 is equivalent to 푛휆 [16].
This can be expressed by Bragg's law:
푛휆 = 2 ⅆ 푠푖푛 휃 (2.21) where 휆 is the X-ray wavelength, ⅆ is the interspacing of the atomic plane, 푛 is an integer, and 2휃 is the angle of the incident beam with respect to the diffracted beam obtained experimentally. 34
Figure 2.8: Illustration of the X-ray diffraction with Bragg's law
Scanning Electron Microscopy (SEM)
SEM, a focused beam of high-energy electrons is used to create microstructure images of material that are derived from electron-material interactions [23]. Typically, a beam of electrons is produced by an electron gun or field emission gun, which is then focused down at the sample surface using optics comprising of electromagnetic coils condenser lens and aperture as demonstrated in Fig. 2.9a.
Upon striking the sample's surface, the focused electron beam gives rise to several different signals like X-ray, auger electron, secondary electrons, backscattered electrons, as illustrated in Fig. 2.9b. The most commonly used are secondary electrons for morphology and 35 topography of samples and backscattered electrons for imaging composition contrast in multiple phases.
Figure 2.9: Description of electron microscopy techniques. (a) Schematic structure of the
electron microscopy column and (b) various signals from the electron-sample interactions
Transmission Electron Microscopy (TEM)
The transmission electron microscope is a very powerful and popular tool for the characterization of materials and a natural complementary technique to PAS [24]. Similar to the
SEM technique described above, the sample is irradiated with a high-energy beam of electrons, and the interactions between the electrons and the atoms can be used to observe features such as the crystal structure and microstructural attributes dislocations and grain boundaries. However, 36 unlike SEM, which creates micrographs by detecting the reflected electrons, the technique of
TEM uses transmitted electrons to create a high-resolution image, as shown in Fig. 2.9b.
Atomic Probe Tomography (APT)
APT is a destructive technique that can provide structural information about atoms' position and chemical identity with high spatial resolution and 3D imaging [25, 26]. The technique's success relies on the sample, which is prepared as a sharp tip that is then biased at a high DC voltage (3-15kV). The combination of low radius (tip size) and high voltage results generates a high electrostatic field at the surface of the tip. The tip of the surface is then subjected to laser or HV pulsing, and atoms are evaporated from the surface and projected upon a position-sensitive detector. Laser or HV pulse acting on the tip of the specimen can be described by a pulse which triggers evaporation of ions, and the time of flight, is when the ion reaches the detector.
However, the technique is often limited by the sample preparation, which requires the preparation of atom probe needles and difficulties with laser-induced atom evaporation, which are required to perform measurements.
Photoluminescence (PL) Spectroscopy
The interaction of light with material often leads to phenomena such as transmission, absorption, reflection, scattering etc. PL is a non-contact and non-destructive spectroscopic method that has been widely used to investigate the optical properties of materials. In PL spectroscopy, light is directed into a material which induces photoexcitation, and electrons are excited from the ground state to permissible excited states. These electrons revert to their 37 original states by releasing a photon which can then be captured to measure the energy between the two states [27].
The schematic of the setup for PL spectroscopy used in this work is presented in Fig.
2.10. The excitation source used is a 456 nm Blue LED light source filtered through a monochromator to remove the tail. The light beam is focused on the sample surface using a fiber optic and a lens. The generated emission spectrum is then collected using a lens and analyzed by a CCD detector. The temperature-dependent photoluminescence (TD-PL) measurements can be performed using a temperature-controlled heating stage where the temperature can be controlled between 83K-673K. The inclusion of TD-PL studies makes it a powerful technique in 38 understanding latent physical and chemical processes which otherwise are undetectable at room temperature, such as luminescence quenching processes.
Figure 2.10: Schematic setup of a PL system
Thermoluminescence (TL) Or Thermal Stimulated Luminescence (TSE) Spectroscopy
The TSE spectroscopy works on the principle that when materials are irradiated by high- energy photons such as UV, x-ray, or γ-rays, electrons get excited from the valence band to the conduction band and generate charge carriers [28, 29]. In the presence of defects, electrons and holes may become trapped, which on subsequent heating can get sufficient thermal energy allowing electron de-trapping and thus allowing the transfer of energy to recombination centers leading to light emission, which can be easily recorded. This is illustrated in Fig. 2.11.
TSE glow curves constructed using emission vs. intensity can also provide very accurate energy levels of traps in the bandgap. Many approaches have been developed to calculate these trap levels energies. In this work, one such approach- the initial rise method. The technique is 39 based on the principle that during the initial rise of the peak, the number of trapped electrons can be assumed to be a constant, the dependence of temperature is negligible. The initial part of the
TSE glow curve peak is then exponentially dependent on temperature according to the equation:
I(T)= constant × exp (-E/kT). However, this assumption is usually valid only if the temperature value is significantly less than the temperature at which peak maxima occurs. This value usually corresponds to the intensity, which is 10-15% of the maximum intensity. 40
Figure 2.11: Schematic demonstrating a simplified TSE process. 41
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CHAPTER III. DEFECT STUDIES IN ION-IRRADIATED IRON
Introduction
Radiation damage is a research topic of high importance, impacting many major areas of science [1-6]. Especially a great deal of research has been conducted on the effect of radiation on materials [7-9]. Continuous exposure to high-energy particles would result in the introduction of many defects in the crystalline lattice of materials. Prior studies involving ion irradiation confirmed the formation of several defect types like interstitials, dislocation loops, vacancy- related defects, and voids in materials [10-12]. These defects can cause detrimental changes in properties, including swelling, hardening, and embrittlement, and are expected to impact how the material reacts to other environments, such as a corrosive medium. Thus, it is critical to understand the nature of these defects and their properties on an atomistic scale. Despite many studies on radiation damage in materials, there is still much that is unknown about how defects are generated at the atomic scale and how those defects interact with the pre-existing microstructure. Much of what is known comes from atomistic simulations, which, while valuable, always have significant uncertainties. Thus, fundamental insights into the damage produced in materials can aid in developing models of damage evolution. While destructive microscopy and imaging techniques such as high-resolution transmission electron microscopy
(HR-TEM) have provided invaluable information regarding irradiation-induced voids and loops, they cannot capture defects on the atomic scale, such as single vacancies or small clusters less than 1.5 nm in size.
Variable energy positron annihilation spectroscopy (VEPAS) can be a reliable tool for its high sensitivity to open volume defects and depth-resolved profiling in the range of few nanometers to several microns and can be used to probe surfaces, thin layers, and near-surface 46 regions in solids. The current study employs the newly developed short-pulsed positron beam in
Dresden, Germany, and presents a detailed quantitative assessment for the depth dependence of the damage profile for ion irradiation, capturing both the size and density of vacancy clusters. It highlights the non-destructive nature of this analysis to describe the atomic-scale damage profile of ion irradiation even in highly defected films.
Experimental Methods
Material Growth And Preparation
Coarse-grained films of pure Fe (labeled as Pure Fe) were fabricated at a temperature of
500 °C on Si (100) substrates in a high vacuum chamber with a base pressure of < 6.6 × 10-6 Pa, using the physical vapor deposition system at the Center for Integrated Nanotechnologies (CINT) at Los Alamos National Laboratory (LANL). The Fe power was operated at 400 watts. The deposition rate was approximately 0.2~0.3 nm/sec, with the total time of the deposition being
4000 seconds. The pristine films were found to have a high density of dislocation and voids.
Ion Irradiation
Ion irradiation of the Fe films was performed with 2 MeV Fe ions at room temperature on the 3 MV Tandem accelerator at the Ion Beam Materials Laboratory at Los Alamos National
Laboratory. To achieve 1 cm in diameter irradiation area, which is needed for PAS characterization, the focused 2 MeV Fe ion beam was rastered using an electrostatic scanner with fixed frequencies of 512 Hz and 64 Hz in horizontal and vertical directions, respectively. The ion fluence was measured using a 4-corner Faraday cup assembly in front of the samples. The target chamber vacuum was maintained below 5x10-8 torr during all irradiations. Irradiations with peak damage of 0.006 to 0.06 dpa were carried out using a fixed damage rate of approximately
2.5x10-4 dpa/s. The beam fluence to damage conversion was based on SRIM Monte Carlo code 47 simulations using the K-P mode with a displacement threshold energy of 40 eV for all the elements involved.
Transmission Electron Microscopy (TEM)
TEM measurements were carried out to investigate the microstructure of both the reference and irradiated samples. Cross-sectional TEM samples were prepared from the focused ion beam (FIB) lift-out method from the deposited Fe films. The lift-outs were created with a
FEI Quanta 3D FEG Dual Beam SEM/FIB with a Ga ion source capable of energies up to 30 kV.
Final thinning was done at 2 keV to reduce the Ga beam damage before examination in the TEM.
After preparation, the TEM samples were imaged using both a JEOL JEM-ARM200CF and FEI
Talos F200X G2 analytical for the unirradiated samples. For the irradiated samples, the JEOL
ARM200CF was used in TEM mode to measure the defect distribution, with images were taken using a Gatan Orius SC200D camera and the Gatan Microscopy Suite v3.3. TEM bright-field
(BF) images were acquired at under focused and over-focused conditions with a defocus value of
2.36 um to observe the cavities present in the samples. The under-focused images were selected to measure the size and density of the cavities. In the under-focused TEM images, cavities appear predominantly as white circles encircled by a darker Fresnel fringe. The cavity size was measured as the inner diameter of the dark fringe surrounding the white dots.
Variable Energy Positron Annihilation Spectroscopy (VEPAS)
VEPAS was performed at the positron facility at the Helmholtz-Zentrum Dresden-
Rossendorf (HZDR) in Dresden, Germany [10-11]. Doppler broadening spectroscopic (DBS) measurements were carried out at room temperature for each sample in the range of 0.05-16 keV using a DC monoenergetic positron beam. A high purity Ge detector was used to detect 511 keV gamma rays produced by positron annihilation at each beam energy with an energy resolution of 48
1.09 ± 0.01 keV at 511 keV. The resulting spectra were characterized by line shape parameters S representing annihilation events with valence (low momentum) electrons [12].
Depth-resolved positron annihilation lifetime measurements were performed in the 0.5-16 keV range at the Monoenergetic Positron Spectroscopy (MePS) beamline facility, also at the
HZDR [11]. The resolution function was determined to be around 250 ps, and positrons were generated with a flux of 106/s. MePS is one of the subsystems of EPOS (ELBE Positron Source) connected to the primary electron beam: Electron Linac for beams with high Brilliance and low
Emittance (ELBE). The positron beam at MePS is created by pair production of ELBE electron beam on a tungsten target, and positrons are generated with a flux of 106/s. At each beam energy, spectra of more than 106 annihilation events were recorded.
A set of Gaussian distribution functions were convoluted with the obtained lifetime spectra to describe the spectrometer resolution function and was confirmed to be around 250 ps.
The obtained spectra were analyzed using the PALSFit software [13] and were decomposed into two lifetime components. An additional third component was found with negligible intensity (<
1%) and lifetime in the ranges of 1~2 ns which can be attributed to the formation of ortho- positronium and thus ignored.
Results And Discussions
Radiation Damage And Positron Stopping Profiles
Figure 3.1a is a schematic illustrating the range of Fe ions in the films, the induced damage profile, and the implantation of positrons. Figure 3.1b displays the radiation damage accumulation profile in terms of displacement per atom (dpa), as generated by the SRIM code when irradiated by 2 MeV Fe ions to a fluence of 5.65 x 1014 ions/cm2. As expected, the damage profile is non-uniform over the irradiation depth, increasing almost linearly from the surface to 49 the peak dpa region, after which it quickly falls off. Two different fluences are considered in this study, corresponding to peak dpa values of 0.006 and 0.06. Figure 3.1c illustrates the intersection of radiation damage profiles with positron stopping profiles at positron implantation energies of
4 keV and 16 keV(26). It is clear that the overlap between the radiation damage curve and positron implantation profile significantly increases with increasing positron incident energy, thus leading to a greater sensitivity of probing damage using positrons at higher average depths.
In fact, at an implantation energy of 16 keV, the overlap compasses the complete range of radiation damage, as shown in Fig. 3.1c.
DBS Measurements
Figure 3.2 shows the DBS measurements S(E) for the pristine film and the two irradiated (0.006 and 0.06 dpa) films. Dashed lines represent the fitting of S(E) profiles obtained using VEPFIT
[14]. The defect parameter S is sensitive to both the size and concentration of defects and is represented on the vertical axis as a function of positron implantation energy (bottom axis) and corresponding mean penetration depth of positrons (top axis) 푥̅. A significant difference in the shape of the S parameter curve between the irradiated and pristine films can be observed in Fig.
3.2. S is higher at shallow depths in the pristine case because of the effect of surface states and positronium formation, then decays with depth and starts to plateau at about 190 nm. S exhibits the opposite behavior in the irradiated materials, increasing with depth due to increased radiation-induced defects at higher depths. 50
Figure 3.1: Damage and positron implantation profiles in the Fe thin films. (a) A schematic
illustrating the probing of induced dpa damage with the positrons. (b) Damage profile as generated from SRIM simulations (Irradiation: 2MeV Fe ions). (c) Overlap of the damage profile in Fe with Makhovian distribution profile of positrons at implantation energies of 4 and 16 keV. 51
Examining these profiles in more detail, S(E) in the pristine film first displays a sharp decrease in the range 0.03-63 nm followed by a gradual decay until it becomes relatively constant. Irradiation to doses of 0.006 and 0.06 dpa lead to a large increase in S as compared to the pristine sample and can be attributed to the formation of small vacancy clusters (as opposed to large voids, which are present in the pristine material and are less effective positron traps) which become the dominant positron traps. Irradiated samples exhibit an increase in the S parameter within the first few keVs of implanted positrons which could be due to back-diffusion of positrons and their annihilation at the surface, as often observed in these materials [15]. The
S(E) profiles shown for the two irradiated samples suggest different defect characteristics with the variation of radiation dose. The sample irradiated at 0.06 dpa exhibits a sharp rise in the S parameter between 0.4-44 nm followed by a modest change until it reaches a plateau at higher implantation depths. In addition, the S values are relatively larger for the 0.06 dpa irradiated sample than the 0.006 dpa sample in the range 0.05-63 nm, indicating more radiation-induced defects have been introduced at the surface region with the increased dose. At greater depths, S is smaller for 0.06 dpa than 0.006 dpa, perhaps due to the growth of some vacancy clusters to larger voids that are not strong positron traps. This will be examined in detail below in conjunction with positron lifetime data.
Calculation of bulk positron diffusion length L+ (to be used in defect density calculations): The bulk positron diffusion length (퐿+푏) within a defect-free material is given by:
퐿+푏 = √퐷+푇푏 (3.1)
Where 퐷+ is the positron diffusion coefficient within the bulk material and 푏 is the positron annihilation lifetime in bulk. They are equal to 1.87 cm2s-1 and 110 ps, respectively, for pure single-crystalline iron. From equation (3.1), 퐿+푏 in iron is estimated to be around 143.42 nm. 52
퐿+ for the three samples was obtained from fitting the S(E) curve using VEPFIT as explained in the results section. The VEPFIT program, introduced by Van Veen et al. in 1991 [14], features a calculation routine based on the solution for the time-averaged positron density equation. The software follows a semi-linear fitting regime where a multi-step procedure is followed by solving a linear minimization problem using the least-squares approach at every step of the non-linear iteration. The S parameters (including bulk, defect, surface, epithermal, etc.) are input as linear parameters that are then used to solve non-linear parameters like diffusion lengths and layer ranges. All the samples were fitted according to “Model 4,” which features calculations for layered structures of defects in materials. The fitting required a two-layered model featuring a surface layer and non-surface region of the film. The fitting determined the width of the surface layer to be less than 40 nm for all the samples and the corresponding 퐿+ values were found to be around 16(1) nm for the reference sample, 11(1) for the 0.006 dpa and 8(0.6) for the 0.06 dpa irradiated samples. These low 퐿+ values in layer-1 are attributed to the high density of defects at the surface. The 퐿+ values corresponding to the non-surface layer of the film are depicted in
Table 3.1 for all the samples. In addition, the reference and 0.006 dpa irradiated samples exhibited relatively good fits (reduced chi-square around 1.7 and 1.5 respectively) as compared to the 0.06 dpa irradiated film (reduced chi-square around 9). In principle, 퐿+ values are discerned from the diffusion of positrons, implanted at low energy, back to the surface, where annihilation occurs at defect-like surface states. From Fig. 3.2, it is clear that the sample irradiated at 0.06 dpa has significantly higher S values at the surface (0.4-44 nm), which implies that the 0.06 dpa irradiated sample should manifest a shorter diffusion length. This implies that the 0.06 dpa sample has a higher defect concentration. However, it is difficult to conclude this with the mean fit values observed for irradiated samples. Thus, using the lower bounds (given by 53 average value minus standard deviation) of the fitted diffusion lengths yields an estimated upper bound on the defect concentrations which is consistent with the PALS data acquired for these samples, where both the fitted defect lifetimes and the average lifetime exhibit higher values in the 0.06 dpa sample than the 0.006 dpa sample, demonstrating higher concentration of defects as expected. This suggests a value closer to the lower bound of 퐿+ better represents the defect concentration. Thus, the positron trapping rates, and defect densities are calculated using the lower limits of fitted diffusion lengths.
Figure 3.2 also shows fits of the S(E) profiles, as obtained from VEPFIT, from which the average positron diffusion lengths (퐿+) of the three samples can be obtained, as summarized in
Table 3.1; 퐿+ in the pristine sample is significantly lower than that calculated for defect-free iron and is expected for a reference film with a high concentration of defects prior to irradiation, consistent with the TEM analysis in Fig.3.5a, which provide higher magnification imaging for quantitative analysis of the initial void density. The uncertainties in the fitted 퐿+ are comparable to those reported in the literature [16]. As to defect densities increase, positron scattering and trapping rates increase, resulting in shorter diffusion lengths. This can be observed here, with the irradiated samples exhibiting smaller 퐿+ values than the reference sample. 퐿+ for the 0.06 dpa irradiated sample features higher uncertainty; however, given that the uncertainty is less than the magnitude, it can still provide reasonable bounds on the defect concentration. In principle, a substantial increase in irradiation dose would significantly change the defect distribution and densities and typically lowers the positron diffusion length. The PAS data also well support this, as when the irradiation dose was increased by a factor of ten from 0.006 to 0.06 dpa, the S(E) profiles showed remarkable differences throughout the depth. Similarly, a sharp increase in the
PALS parameters, including individual positron lifetime components and average lifetime 54 values, was observed. These experimental observations and trends are well reflected when lower bounds of the positron diffusion lengths are taken. This also allows the estimation of an upper bound on the defect concentration and reduces the magnitude of uncertainties associated with calculating the average values of defect densities due to the large spread of diffusion length at higher dpa dose.
Table 3.1: Table of 퐿+ values (average positron diffusion length) and the two lifetime
components extracted for each sample.
Sample 푳+(Mean, nm) 푳+(Min, nm)
Reference 107 ± 15 92
Irradiated (0.006 dpa) 70 ± 14 56
Irradiated (0.06 dpa) 67 ± 33 34
Sample 흉 (Range, ps) 흉 (Range, ps) ퟏ ퟐ
Reference 153- 182 355- 373
Irradiated (0.006 dpa) 207- 258 376- 422
Irradiated (0.06 dpa) 230- 279 382- 434 55
Figure 3.2: Doppler broadening measurements for the reference and irradiated samples featuring defect parameter S as a function of implantation energy and depth. Dashed lines represent fits of
S(E) profiles by VEPFIT. The inset shows calculations of S from the 511 keV peak at every
energy.
Depth-resolved PALS Measurements
PALS has been established as the key tool to determine the size and nature of open volume and vacancy-related defects [12]. However, its application to ion irradiated materials was not possible because of the absence of pulsed positron beams. Here a monoenergetic pulsed positron beam with controlled implantation energy is used to probe the sub-surface region of the ion-irradiated films with a penetration depth from a few to hundreds of nanometers. At each depth, a positron lifetime spectrum (Fig. 3.3a shows a representative lifetime spectrum with 2 components) is collected and analyzed, as explained in the experimental methods section. 56
Positrons at each depth annihilate with electrons either from a delocalized state in bulk or from a trapped state in a vacancy cluster. When no trap is present, the positrons annihilate in bulk with only one lifetime equal to the bulk lifetime (휏b), which is well known for most materials [17].
When vacancy clusters are present, they efficiently trap positrons, delaying annihilation and leading to longer lifetimes because of the lower electron density inside vacancies. It is worth noting that positrons are typically sensitive to trapping at vacancy centers having concentrations as low as 0.1 ppm [12]. Fitting the lifetime spectrum often gives two or three lifetime components with their intensities depending on the number of traps and the feasibility of resolving their lifetimes. The lifetime value and intensity are related to the size and density of vacancy clusters, respectively. The lifetime values for vacancy clusters in Fe up to 15 vacancies are known from density functional theory calculations [17] and will be used here for data interpretation. It should be mentioned here that grain boundaries should not affect positron trapping or diffusion in large-grained materials as are present in the current study.
Fitting the spectra of all samples gave two components (휏1 and 휏2, Fig. 3.3b) and their intensities (퐼1 and I2, Fig. 3.3c) for each sample; Figure 3.3d shows the average lifetime value.
More details about the fitting and applying of the two-defect trapping model are discussed in the next section in conjunction with defect density calculations. The short-lived components (휏1) in all samples exhibit lifetimes higher than the previously reported bulk value 휏b in the range of
106-110 ps for Fe [18], the upper limit is considered in this study. This indicates that all positrons were trapped at defects and that no positron annihilation occurs in bulk even in the reference sample, another indicator of the high density of defects in the starting material, as shown in Fig. 3.5a later. Accordingly, the two fitted lifetimes 휏1 and 휏2 represent two defect groups. The observed increase in the lifetime values when the material is subjected to radiation, 57 in the order 0.06 dpa > 0.006 dpa > reference, indicates the growth in the size of vacancy clusters with increasing dose. In the reference film, 휏1 is within the 153-182 ps range, which is higher than 휏b (110 ps) and lower than the monovacancy lifetime 휏1V (175-190 ps) for Fe [17], the upper limit is considered in this study. Such lifetime values are known to be associated with dislocations [17, 19]. The second lifetime 2 in the pristine film is around 350 ps, indicating the presence of large clusters containing about 10 vacancies (10V). In the irradiated samples, the first lifetime component 휏1 exhibits values ranging from 207-258 ps, indicating the presence of divacancies (2V) in the lower dose sample that grows to small vacancy clusters pertaining to the size of 4 vacancies (4V) when the dose is increased to 0.06 dpa. A similar trend is observed for the second lifetime 휏2 where the sample irradiated to 0.06 dpa has higher long-lived lifetimes
(ranging from 382-434 ps) as compared to the lower dose sample (376-422 ps). These values indicate the formation of clusters containing more than 15 vacancies (15V) [17]. 58
Figure 3.3: Summary of depth-resolved PALS experiments. (a) A visual representation illustrating how a positron lifetime are decomposed to separate lifetime components 휏1 & 휏2. (b)
Lifetime components 휏1 & 휏2. corresponding to the two different defect groups (c)
Corresponding intensity components I1 and I2. (d) Average positron lifetime values, the weighted
average of the two lifetime components, suggesting that the 0.06 dpa irradiated sample has
higher defect content, as expected. The grey area marked at 2 keV indicates values obtained at
the surface region. 59
Defect Density Calculations
Standard trapping models can be applied to calculate the defect densities where the trapping rate (퐾) is proportional to the defect concentration (C), where , a defect-specific positron trapping coefficient gives the proportionality constant
퐾 = C (3.2)
The standard two-defect trapping model [12] assumes two defect groups and is often applied to decompose lifetime spectra into a sum of three exponentials. The first is a bulk-like component with reduced lifetime (휏0) compared to the bulk defect-free lattice (휏b) because the presence of defects scatter positrons and limit their diffusion and lifetimes. The other two components exhibit extended lifetimes corresponding to two different positron trapping centers.
However, when large concentrations of defects are present in the samples, nearly all the positrons become trapped at defects, and no annihilation signal arises from the bulk. In this case, 휏0 becomes negligible, and only two lifetime components are observed, both of which can be treated as defect lifetimes. As explained above, the samples used in this work contain a large concentration of defects with only two lifetime values sufficiently higher than that of the bulk lattice; thus, since 휏0 was not observed in the measurements, the two components 휏1 & 휏2 are treated as defect-specific lifetimes with relative intensities I1 and I2 corresponding to two groups of defects. As per the lifetime data shown in Fig. 3.3b, the first group of defects (D1) can be assigned to relatively shallow positron traps (associated with 153-182 ps lifetimes) such as dislocations or small vacancy clusters (associated with lifetimes greater than 190 ps) while the second group of defects (D2) is assigned to large vacancy clusters extending to 15V [17]. There is a significant increase in the lifetime, i.e., in the cluster size, around 190 nm for the higher dose, resulting from the depth distribution of the 2 MeV Fe ions used for irradiation. 60
Since these films have saturated positron trapping at defects, the standard method for calculating trapping rates from PALS is not applicable. Instead, we calculate the trapping rate from L+ and combine both variable energy DBS and PALS to yield an effective method for accurately assessing depth-resolved defect densities even in highly defected films.
It is well-known that positron diffusion lengths derived from back diffusion of implanted positrons can still provide an accurate estimation of defect density even in the case of saturation trapping [21, 22]. For this analysis, the total trapping rate of defects in each film can thus be given by (3.3)
2 1 퐿+푏 퐾 = ( 2 − 1). (3.3) 푇푏 퐿+
The value of the trapping rate obtained from equation (3.3) is an average value containing the combined contributions of defects present in both groups D1 and D2 measured by PALS. It, therefore, can be decomposed into individual components by using values of the intensities obtained from positron lifetime experiments [22] as given by the following equation:
퐾푖 = 퐼푖퐾 (3.4)
Where 퐾푖 (푖= 1, 2) are the individual trapping rates corresponding to the intensity component from each defect group at a given implantation depth and are represented respectively in Fig.
3.4a. Note that K is the total trapping rate calculated from Eq. (3.3) for each sample.
As discussed above, the concentration of defects can then be directly obtained from Eq.
(3.2) by using the defect-specific positron trapping coefficient as the constant of proportionality, assuming it to be relatively constant within one defect group. In the literature, the typical specific trapping coefficient value used for the Fe monovacancy is 1.3 x 10-14 m3s-1 [23, 24], while for large vacancy clusters in Cu, a value of in the order of 5 x 10-13 m3s-1 was reported [25], which has been successfully applied in previous reports [23] for calculations of vacancy clusters in iron 61 and thus will be applied here for the calculation of defect density. For the reference sample, where the lifetime values obtained in group D1 are less than that of monovacancies, the value for the specific trapping coefficient used is for Fe dislocation loops, assumed to be 7 x 10-5 m2/s, as previously reported [25]. Finally, it should be mentioned that the accuracy of the defect density calculations is limited by the un-availability of positron trapping coefficients for all cluster sizes.
They have only been reported for single vacancies, dislocations, and large clusters.
Figure 3.4a displays the trapping rates as a function of depth. The calculated defect densities corresponding to the two defect groups are provided in Figs. 3.4b, 3.4c, and 3.4d. The dislocation density in the pristine film calculated from PAS measurements using this analysis
(presented in Fig. 3.4b) is in very good agreement with the TEM measurements in Fig. 3.5.
In accordance with the lifetime data provided above, the defect densities corresponding to both defect groups comprising of small-sized point defects (Fig. 3.4c) as well as large clusters (Fig.
3.4d), increases significantly when the material is irradiated to 0.006 dpa and displayed a further increase when the dose was elevated one order of magnitude to 0.06 dpa. The density of small clusters increases with depth while it decreases for larger clusters at high depths. This is also reflected in the depth dependence of the ratio between the density of large clusters and small clusters (Fig. 3.4e). The decrease in the density of group D2 defects (which represents large clusters >15 vacancies) with depth (Fig. 3.4d) is associated with an increase in their size, as indicated from the increase in lifetime in Fig. 3.3b, which is a direct result of the increase in cluster size. This implies that these defects have agglomerated to form larger clusters and reveal that larger clusters are stable at higher penetration depths away from the surface. No monovacancies were detected as they are often highly mobile unless trapped at impurities, which 62 is not the case in the current measurements, possibly because of lower impurity concentrations or the high density of sinks in the films.
Figure 3.4: The properties of ion-beam-induced defects in the thin films from PAS. It features:
(a) trapping rates obtained from positron diffusion lengths and Eq. (3.3) and (b) dislocation
-2 -3 densities (m ), D1, calculated for the pristine sample. (c) Densities of small clusters (m ), D1.
The inset shows a schematic of the types of small clusters. (d) Densities of larger clusters (m-3),
D2. The inset shows a schematic of the types of larger clusters. (e) Ratio of defect densities 63
calculated for irradiated samples corresponding to the two defect groups D1 and D2. (f) Overall
defect density obtained for irradiated samples by summation of individual defect densities
corresponding to defect groups D1 and D1. All samples exhibited very high defect
concentrations.
TEM Measurements And Calculation Of Cavity And Void Density
TEM images of the pristine film presented in Fig. 3.5 show 0.9 m thick films with high densities of dislocations and voids. These micrographs reveal that the pristine films, deposited on a Si substrate, contain weakly-textured polycrystalline grains about 230 nm in diameter. The dislocation density is estimated to be about 1.5 x 1014 /m2. There is also a high degree of porosity/voids in the films. These features will modify the behavior of positrons in the material, complicating the analysis of the positron spectra. Moreover, they will also modify the propagation of cascades and defects formed during irradiation. 64
Figure 3.5: TEM images of the pristine sample. (a) Thickness of the film grown on a Si substrate with an average grain size of ~ 230 nm, (b) Dislocation density obtained in the highlighted cross-
section average grain size, and (c) Presence of voids throughout the film. The high densities of
dislocations and voids indicate that the pristine films are highly defective.
The microstructures of the higher dose irradiated films were directly observed with TEM bright-field (BF) imaging following irradiation (Fig. 3.6). The cavity density and average cavity size for the pristine sample and the sample irradiated to 0.06 dpa are presented in Figs. 5a and
5b. The cavity density is on the order of 1x1023 m-3 for the two samples. That is, the density did not change significantly upon irradiation. For the reference case, the average cavity size is from
1.8 to 2.0 nm and shows a dependence on depth, while it ranges from 1.6 to 2.0 nm for the sample irradiated to 0.06 dpa and shows a stronger dependence on depth. The cavity size for the irradiated sample was measured to be 1.94 nm at a depth of 130 nm and decreased to 1.61 nm at a depth of 340 nm. 65
Figure 3.6: Post-irradiation defect quantification from TEM micrographs of the irradiated
sample. (a) cavity density and (b) average cavity size plotted as a function of depth from the surface. Figures to the left of the plots show the micrographs used for the cavity calculations for
pristine and 0.06 dpa samples.
The TEM observations show that the cavity density, already very high in the pristine film, is unchanged to a dose of 0.06 dpa. However, the average cavity size decreased after irradiation. At the same time, the PAS results showed an increase in the density and size of smaller defect clusters with irradiation. This suggests irradiation is shrinking the larger-scale
TEM-visible cavities present in the samples while simultaneously creating and growing the smaller scale vacancy clusters that PAS can observe. The data here reveal that void shrinking can happen at least for irradiation dose below 0.1 dpa. 66
Further studies are needed to illustrate if this behavior would last at higher doses. The interpretation above is supported by recent work by Fellman et al. [26], who simulated voids and other vacancy type defects interacting with cascade events induced by irradiation in BCC tungsten at room temperature. They observed that cascades overlapping with preexisting voids resulted in a significant decrease in the original void size. In these simulations, not only did preexisting voids decrease in size, but new vacancy clusters survived after cascade overlap. Our combined PAS and TEM results demonstrate for the first time a very similar behavior on highly defective Fe films, direct experimental validation of the modeling results. Because of the high initial cavity density in the reference films in our study, the probability of cascade overlapping with pre-existing cavities is high, leading to a reduction in the void size and a corresponding increase in the surviving vacancies after cascade events, due to the sink behavior of the voids with cascade-produced interstitials. This mechanism is schematically illustrated in Fig. 3.7. The surviving vacancies form smaller vacancy clusters defects that are too small for TEM to image 67 but are detected by PAS. Thus, by combining PAS with TEM on ion-irradiated materials, new insight into the evolution of radiation damage has been revealed.
Figure 3.7: The schematic shows the pre-existing voids in the reference films and the overlap
between them, and the cascade induced by irradiation.
Conclusions
This work shows the effectiveness of PAS in quantifying the size distribution and concentrations of vacancy clusters associated with ion irradiation. The recent development of pulsed positron beams with good timing resolution provides a new capability for applying PALS on ion-irradiated samples. Future studies of defect formation in magnetic materials such as iron could also greatly benefit from the development of spin-polarized positron beams. 68
PALS, combined with more traditional TEM observations, reveals an important new mechanism for cascade-void interaction where cascade formation leads to shrinkage in the size of preexisting voids associated with an enhancement in the formation of smaller vacancy clusters. These results provide direct evidence of the ability of internal surfaces to act as sinks for defects, providing at least a temporary reduction in the growth of larger-scale defects responsible for material failure. They also demonstrate the ability to non-destructively probe the damage produced during ion irradiation, opening the way for new types of radiation damage studies and more direct validation of modeling and simulations.
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CHAPTER IV. DEFECT STUDIES IN ION-IRRADIATED FE-CR ALLOY
Introduction
Chromium alloyed ferritic/martensitic steels are leading candidates for several advanced reactor applications, including the development of fusion reactors, accelerator-driven systems, and Gen IV reactors [1, 2]. In particular, alloys such as HT-9 (12Cr), T-91 (9Cr), or F82H (8Cr) are being considered for these applications. These alloys are attractive since heat treatments can tailor their properties, and their structure – the body-centered crystal structure of the alloys and the martensitic microstructure – enhance their radiation tolerance of the materials, leading to a low swelling rate. Ultimately, the material response to radiation damage is governed by point defects; their evolution leads to larger-scale microstructural changes (e.g., void formation) and materials degradation on a macro scale, including swelling and embrittlement. The chromium content of these alloys is, of course, tied to their corrosion resistance. Higher chromium content leads to higher corrosion resistance in oxidizing environments while also stabilizing the bcc phase and, in conjunction with carbon, is responsible for the desirable mechanical properties [3,
4].
A relatively common notion in the literature is that a Cr concentration around ~8-10%
(wt. %) is a safe compromise to achieve a balance of different properties necessary for the aforementioned applications [5]. It is known that the addition of even small quantities of Cr remarkably reduces radiation-induced swelling, an effect that becomes more prominent with increasing Cr alloying up to ~8-10 wt. %, after which it may or may not lead to further improvements, depending on the irradiation conditions [6-8]. However, the reason behind this behavior is not yet well understood, and the hypothesis that the formation of Cr- rich ’ phase precipitates act as additional recombination centers for drifting point defects cannot be supported 73 for lower Cr content materials since no ’ has been observed in 9%-Cr materials, though, higher
Cr content materials may show ’ formation [9, 10]. Recently, in a comparative study [11], ’ was shown to form after neutron-irradiation but not ion-irradiation in Fe-12%Cr.
A relatively accepted explanation entails the interaction of Cr atoms with self-interstitial atom (SIA) clusters, which act as sinks for point defects [12]. Interestingly, most of these investigations were based on treatments with higher doses of irradiations, and there is not much information on the mechanism in low dose irradiation regimes, where the interactions are, in some sense, in their simplest form. This knowledge gap has been a long-standing problem in the field. It necessitates the need for a fundamental study investigating the role of chromium alloying on the manifestation and evolution of defects in these materials at the atomic level after low dose irradiation. This requires an atomistic probe with high sensitivity to point defects, and, so far, positron annihilation spectroscopy (PAS) is the only known reliable technique capable of performing such measurements [13-15].
Here, by performing the first depth-resolved PALS study on ion-irradiated Fe-Cr alloys and combining PALS with DBS, we extract the density of small and large clusters as a function of depth in both pristine and irradiated films and examine the change of the size of radiation- induced clusters with Cr content. The most important findings of this work come from the PALS measurements, which reveal that Cr in pristine films stabilizes a complex defect center involving
Cr and vacancy clusters that act as defect sinks upon exposure to irradiation. It also shows that a higher percentage of Cr (18%) leads to highly non-uniform defect distribution throughout the depth of the irradiated region, despite no evidence of ’ was found before or after ion-irradiation.
The study was carried on Fe, Fe-Cr alloyed films to control Cr percentage throughout the growth process better. 74
Experimental Methods
Material Growth And Preparation
Coarse-grained films of pure Fe (labeled non-alloyed) and FeCr alloys with varying Cr concentrations of 8 and 18 wt. % (labeled Fe-8Cr and Fe-18Cr) were fabricated at a temperature of 500 °C on Si (100) substrates in a high vacuum chamber with a base pressure of < 6.6 × 10-6
Pa, using the physical vapor deposition system at the Center for Integrated Nanotechnologies
(CINT) at Los Alamos National Laboratory (LANL). For Cr-alloyed film deposition, Fe and Cr sources were used simultaneously with independent shutters and power supplies. The Fe power was operated at 400 watts and the Cr power at 40 watts to obtain the chemical composition of
Fe-8Cr and 70 watts to obtain Fe-18Cr. The deposition rate was approximately 0.2~0.3 nm/sec, with the total time of the deposition being 4000 seconds.
Ion Irradiation
Irradiation of the non-alloyed and Cr-alloyed films was done with 2 MeV Fe ions at room temperature, using the 3 MV NEC Pelletron tandem accelerator at the Ion Beam Materials
Laboratory at LANL. The diameter of the irradiation area was 1 cm which was achieved via rastering of the 2 MeV focused Fe ion beam through an electrostatic scanner with fixed frequencies of 512 Hz and 64 Hz in the horizontal and vertical directions. A radiation damage level of around ~0.06 dpa was attained by using a damage rate of 2.5 x 10-4 dpa/s. The beam fluence-to-damage conversion is based on SRIM Monte Carlo code simulations using the K-P mode with a displacement threshold energy of 40 eV for all the elements involved.
Variable Energy Positron Annihilation Spectroscopy (VEPAS)
VEPAS measurements were performed at the positron facility at the Helmholtz-Zentrum
Dresden-Rossendorf (HZDR) in Dresden, Germany [16]. Doppler broadening spectroscopy 75
(DBS) was performed using a continuous monoenergetic beam, capable of positron implantations in the range of 0.04-17 keV [17, 18]. A high purity Ge detector, with an energy resolution of 1.09 ± 0.01 keV at 511 keV, was used to record the characteristic 511 keV gammas at each beam energy. The experimental data were interpreted by measuring the energy deviation in the characteristic 511 keV annihilation gammas induced by the Doppler shift of the low momentum valence electrons and were represented in terms of the line-shape defect parameter known as the S parameter. Details about DBS spectroscopic measurements and analysis can be found elsewhere [19, 20].
Depth-resolved PALS spectroscopic measurements were carried out at the
Monoenergetic Positron Spectroscopy (MePS) beamline facility at the Electron Linac for beams with high Brilliance and low Emittance (ELBE) positron facility, also at the HZDR [16]. The measurements were recorded at beam energies between 2-16 keV, with a positron flux of approximately 106/s. At each beam energy, spectra of more than 106 annihilation events were recorded. A set of Gaussian distribution functions were convoluted with the obtained lifetime spectra to describe the spectrometer resolution function and was confirmed to be around 250 ps.
The obtained spectra were analyzed using the PALSFit software [21] and were decomposed into two lifetime components. An additional third component was found with negligible intensity (<
1%) and lifetime in the ranges of 1~2 ns which can be attributed to the formation of ortho- positronium and thus ignored.
Atomic Probe Tomography (APT)
APT measurements were carried out using a CAMECA local electrode atom probe
(LEAP) 4000 X-HR system in the laser pulse mode via a 355 nm wavelength picosecond laser with a laser energy per pulse of 70 pJ, a pulse repetition rate of 250 kHz, and a detection rate of 76
0.002-0.005 ions per pulse. The base temperature and pressure were maintained at 45 K and <2 ×
10–11 Torr. The Integrated Visualization and Analysis Software (IVAS 3.8.0), also provided by
CAMECA was then used to reconstruct and analyze measured data.
Results And Discussions
Radiation Damage And Positron Stopping Profiles
Figure 4.1 shows depth-resolved radiation damage profiles for non-alloyed and binary
Fe-Cr alloyed films with different Cr compositions (8 and 18 wt. %) when irradiated by 2 MeV
Fe ions with a fluence of 5.65x1014 ions/cm2. The profiles in Fig. 4.1 show non-uniform damage accumulation as a function of depth, with damage gradually increasing until it reaches a peak value around 600 nm, followed by a rapid decay. To estimate the initial depth distribution of incident positrons, stopping profiles are calculated using the Makhovian distribution function
[22] and are plotted at energies of 4, 8, 12, and 16 keV for comparison. As evident from the positron stopping profiles in Fig. 4.1, the mean range for positron probing at lower positron incident energies is quite short but provides a relatively good depth-resolved resolution. With the increase in the incident positron energy, the peak of the positron stopping profile curves becomes broader, resulting in a relatively large region of positron probing but with a decline in the resolution for the depth-resolved measurements. The plots in Fig. 4.1 show a good overlap between the dpa profile and positron stopping profile at 16 keV and indicate the possibility of probing the entire irradiated region via PAS. 77
Figure 4.1: Damage and positron stopping profiles for Fe and FeCr alloyed films. There is a
significant overlap of the positron stopping profiles in the range of 4-16 keV with the damage
profile created within Fe and FeCr alloys when irradiated with 2MeV Fe ions (as determined
from SRIM calculations).
Depth-resolved Doppler broadening Spectroscopy (DBS) measurements
S(E) profiles are plotted in Fig. 4.2a as a function of positron implantation energy and depth obtained from DBS measurements. The S parameter, which is a measure of the fraction of positrons annihilating with low-momentum valence electrons, is directly proportional to the overall defect content in the material. The measurements show differences in the S(E) characteristics with the variation of Cr content. Relatively high S(E) values within the first couple of nanometers of the surface were observed for all pristine films and could be due to forming a small passive oxidation layer atop the film frequently observed for these materials. 78
Furthermore, when probing from the surface to the near-surface region, a decay in the S(E) values with increased depth was observed for the pristine films and showed notable trends for each sample. The Fe-8Cr alloy exhibited a sharp drop in the S parameter within the first few nanometers of the surface, after which it remained relatively constant to higher depths. In contrast, the decay of the S parameter is more gradual for Fe-18Cr alloy and does not seem to be approaching a plateau until a depth of 133 nm. This indicates a continuous presence of large clusters in the pristine Fe-18Cr film, extending more than 100 nm deep into the sample. After irradiation, all films exhibit an enormous increase in S values throughout the whole depth of the film, though S is significantly higher for the pure Fe film. The figure also shows that S values are lower at all depths for the Fe-8Cr sample than the Fe-18Cr sample, indicating that the Fe-8Cr alloy has less overall defect content after irradiation and thus higher radiation tolerance. To further illuminate the effect of radiation on the transformation of defect structures, the relative change in the S parameter (Srel) is calculated as the fractional change in S parameter values before and after irradiation and shown in Fig. 4.2b. For depths below the surface region (>14 nm), the high Cr composition alloy (Fe-18Cr) sample exhibited relatively more significant changes in S and strong depth dependence as compared to the Fe-8Cr alloy and the non-alloyed film. In fact, for the Fe-8Cr sample, the difference between S before and after irradiation seems constant throughout the depth, an indication that radiation-induced changes in defect characteristics are uniform throughout the irradiated region, which is not the case for Fe-18 Cr alloy.
The calculated positron diffusion lengths (L+) from the fitting of S(E) profiles by
VEPFIT are summarized in Fig. 4.2c and will be used in defect calculations. VEPFIT is a standard program in the PAS field [23]. It features layered fitting of the depth-dependent S 79 parameters based on the solution for the time-averaged positron density equation. The details about the program, layered fitting of S(E) profiles, and the calculations of the L+ values for pure
Fe films were recently reported [24]. Similar calculations and fitting procedures were applied here to determine the defect parameters in the present alloyed films assuming two layers: a highly defective thin layer on the surface due to possible oxidation followed by the sub-surface and bulk of the film. The thickness of the surface layer was determined to be < 20 nm for Fe-8Cr and <30 nm for Fe-18Cr alloy films, respectively. The positron diffusion lengths in this region were significantly short (<10 nm) for all films and will not be discussed further. For the region representing the sub-surface and bulk of the film, the L+ values for all the samples were less than that in defect-free bulk (~143 nm) [24]. The shorter diffusion lengths (< 100nm) confirm the presence of defects in the pristine films. The alloyed samples were found to have significantly higher L+ values than the non-alloyed sample after irradiation and are in complete agreement with the observed S(E) profiles. 80
Figure 4.2: DBS measurements showing depth-resolved S profiles and diffusion lengths. (a) S(E)
profiles displaying S parameter values at each positron implantation energy and the
corresponding depths sampled within the material. (b) Relative change in S parameter of irradiated films as compared to the pristine films. (c) Positron diffusion lengths calculated from
the fitting of S(E) profiles using VEPfit.
Depth-resolved Positron Annihilation Lifetime Spectroscopy (PALS)
Fig. 4.3 represents the first depth-resolved PALS measurements on ion-irradiated Fe-Cr alloys. The measurements revealed two long-lived components, both higher than the positron lifetime in defect-free iron (110 ps) [25]. The absence of a defect-free lifetime component in these measurements indicates a high concentration of defects in the films, as discussed in the previous chapter on irradiated Fe. Figure 4.3a shows the lifetime component (휏1) corresponding 81 to defects consisting of dislocations and small vacancy clusters. The alloyed pristine films exhibited higher lifetime values of 160-195 ps versus 145-160 for the pure Fe pristine film.
Previous work has reported positron lifetimes of 140-160 ps for dislocations in Fe-Cr alloys [26].
However, the lifetime measured here is much larger (160-195 ps), and it is most likely a combination of lifetimes from dislocations and mono- and di-vacancies, which are difficult to individually resolve with the current timing resolution. Thus, these measurements indicate the presence of small-sized vacancy defects in the pristine alloyed films, which is consistent with what has been established before: that the addition of Cr stabilizes di-vacancies up to room temperature [27]. In fact, the pristine alloyed films exhibited an appreciable increase in the intensity I1 (Fig. 4.3b). Again, this increase is likely due to the presence of mono and di- vacancies, which are more effective traps than dislocations, the dominant positron traps responsible for the short lifetime components in the pristine non-alloyed films. Thus, the increased positron lifetimes in the pristine Fe-Cr films indicate the presence of vacancy complexes being stabilized by incorporating chromium. A previous study showed that the addition of Cr reduces the formation of defects with minima around 8-12 wt.% Cr [12]. Figure
4.3c shows the defect lifetimes, 휏2, corresponding to large vacancy clusters, and their corresponding intensities, I2, are plotted in Fig. 4.3d. For comparison, the calculated lifetime values for cluster sizes of 10 vacancies (휏10푉) and 15 vacancies (휏15푉) [25] are given in Fig. 4.3c.
It is clear that alloying with Cr led to the formation of larger-sized clusters in the pristine samples prior to irradiation as the lifetimes are larger with increasing Cr concentration, Fe-18Cr
> Fe-8Cr > Fe. In addition, the size of clusters also appears to be dependent on depth, as evident from the increase in lifetimes for the positron implantations at higher depths. This is well 82 supported by the sharp drop in corresponding intensities (Fig. 4.3d), which is expected since, as more of the smaller vacancy clusters agglomerate to form larger clusters, their number reduces.
After irradiation, 휏1 increased in all films. In pure Fe film, it encountered a large increase of more than 250 ps, indicating the formation of vacancy clusters of 4 vacancies. I1, however, was relatively unchanged before and after irradiation. This large increase in 휏1 without a corresponding change in I1 with irradiation in pure Fe illustrates that irradiation creates small vacancy clusters and positrons switch from being trapped at dislocations to being trapped at small vacancy clusters. This is because vacancy clusters are stronger traps for positron with high binding energies of several eVs. In the alloyed films, a relatively smaller increase in 1 is accompanied by a decrease in its intensity I1 was observed. This indicates that the large concentration of small clusters pre-existing in the alloyed films acts as sinks for interstitials and vacancies, thus exhibiting an increase in size reflected in the increase of 휏1 and decrease in their intensity I1. 휏2 significantly increased after irradiation for pure Fe and Fe-Cr samples, signaling the increase in large vacancy clusters up to 50 vacancies. However, no significant increase was observed for Fe-18Cr. It can also be seen that pure Fe a relatively greater extent of positron trapping at larger clusters, as evident from the increase in I2. This signals a higher concentration of large vacancy clusters and voids in pure Fe films. 83
Figure 4.3: Defect lifetimes and intensities as determined from depth-resolved PALS
measurements. (a) Short-lived positron lifetimes 휏1 belonging to dislocations and small-sized
vacancy defects and (b) the corresponding intensities I1 associated with 휏1. (c) Long-lived
positron lifetimes 휏2 belonging to large vacancy clusters and voids and (d) their corresponding
intensities I2. The solid horizontal lines in (a) and (c) indicate the established lifetimes for
specific vacancy clusters.
Depth-resolved Defect Density Estimation
An important goal of this study is to estimate the depth-dependent density of vacancy clusters in the films as a function of irradiation and chemistry. As noted above, previous studies featuring defect quantification in ion-irradiated materials greatly relied on TEM measurements 84 and were limited to the evaluation of large-sized voids and cavities (greater than 1-2 nanometers). Likewise, past PAS studies, though focused on atomistic scale defects, were limited to bulk measurements due to the lack of pulsed positron beams. Here, this work reveals depth- resolved defect densities for ion-irradiated Cr-alloys by combining the depth-resolved PAS measurements, DBS and PALS, and demonstrates the effect of Cr alloying on defect concentrations throughout the irradiated depths. However, the depth resolution is limited by the positron implantation profile and the positron diffusion length.
Nevertheless, it is a good indicator for the change of defect density with irradiation depth.
The calculation method is discussed in detail in the previous chapter. First, the L+ values obtained from the VEPFIT fitting of S(E) profiles in Fig. 4.2 are used to estimate the total trapping rate (퐾) for each film. The total trapping rate 퐾, which is an overall contribution from both small-sized vacancy defects as well as large vacancy clusters, is further reduced to individual trapping rates (퐾푖) as a function of depth using the depth-resolved intensity components, I1 and I 2, from the PALS measurements in Fig. 4.3. Finally, the corresponding defect densities (C) are calculated by dividing the individual trapping rates by a suitable specific positron trapping coefficient (µ) as described below in Eq. (4.1):
퐶 = 퐾/µ (4.1)
It is important to note here that the accuracy of the above estimation of the defect concentration is limited by the availability of µ for specific defect sizes. The typical value of µ
-14 3 -1 used for a Fe monovacancy (µ1V) found in the literature is around 1.3 x 10 m s [28, 29]. It is also known that the value of the specific trapping coefficient is linearly related to size for larger vacancy sizes (~ up to 8) and, thus, can be obtained by multiplying the number of vacancies with
µ1V [30]. Therefore, to calculate defect densities for small-sized vacancy defects, the µ values are 85
obtained by multiplying µ1V by average vacancy sizes as represented by the lifetime bands in
Fig. 4.3a. For large-sized clusters, µ ~ 5x10-13 m3s-1 has been reported for Cu [29, 31], and has been used previously to estimate defects in Fe [29]. The same has been used here without scaling with the number of vacancies as it is accepted in the field that the positron trapping coefficient saturates around 15 vacancies. Thus, the accuracy in calculating the density of large clusters is tied to the correctness of this assumption. The standard trapping model and the calculation method of trapping rates and defect densities have been described in detail elsewhere [15, 24].
Figure 4.4 summarizes the calculated trapping rates (Fig. 4.4a and 4.4b) and concentrations of small and large clusters for both alloyed and non-alloyed samples (Fig. 4.4c and 4.4d). As discussed above, it is impossible to resolve the 휏1 lifetimes of the pristine samples further into individual defects like dislocations and other small vacancy defects. Thus, an accurate estimation of their defect densities is not possible and thus not provided. However, upon irradiation, positrons stop being trapped at dislocations; instead, they become trapped at small- sized vacancy defects because of their higher binding energies for positrons, as mentioned above.
Figure 4.4c presents the defect concentrations for alloyed and non-alloyed films after irradiation, and it is clear that the addition of Cr significantly reduces the density of the small-vacancy clusters in the irradiated alloyed films. Congruent with the trends observed for small-sized point defects, a remarkable reduction in the density of large clusters and voids, more than an order of magnitude, was observed for the irradiated alloyed films as compared to the non-alloyed film
(Fig. 4.4d). It is important to note here that the lifetime intensities I 1 and I 2 in Fig. 4.3 do not directly reflect the density of defects; rather, they indicate the relative change between the density of small and large vacancy clusters. This is because of the high concentration of defects that lead to saturation of positron trapping in the samples (i.e., no bulk annihilation because of 86 the high density of defects in these samples). However, despite this saturation, the use of the positron diffusion length, which is not affected by saturation of positron trapping, allowed us to estimate the total trapping rate, then scale it to the small and large vacancy clusters throughout the depth based on lifetime intensities I 1 and I 2 to estimate the defect densities. Fig. 4.4e compares the density of large vacancy clusters before and after irradiation for the three different chemistries, demonstrating the strong effect of Cr alloying on reducing the level of large vacancy clusters induced by irradiation. We note that the “large” vacancy clusters discussed here that are observed to increase in the pure Fe film after irradiation are not TEM-visible voids that are observed to remain constant in density after these doses [24].
The results from the combined PALS and DBS studies of these ion-irradiated films highlight the role that Cr has in modifying the evolution of defects in Fe-Cr alloys. Cr leads to the stabilization of small vacancy clusters within the material in the pristine films, as revealed in Figs. 4.3a and
4.3b. This suggests that Cr solutes are binding to vacancies, increasing the background density of vacancies in the material. Upon irradiation, Cr clearly suppresses the density of large vacancy clusters (Fig. 4.4d). This suggests a smaller bias for interstitial annihilation that the sinks for interstitial annihilation are more distributed. This enhances the mutual annihilation of interstitials and vacancies and reduces the concentration of large vacancy clusters and voids. This is similar to the sink effect exhibited by oxide precipitates in nanostructured ferritic alloys, which trap He and prevent the formation of large gas bubbles. However, we should note that the high level of defects and porosity in the pristine films may impact the interaction of radiation-induced defects with Cr-vacancy clusters and their evolution. 87
Figure 4.4: Estimation of trapping rates and defect densities in alloyed samples. (a) Trapping rate
for small-sized vacancy defects in irradiated films. (b) Trapping rate for large clusters in films
both before and after irradiation. (c) The concentration of small-sized vacancy defects in 88
irradiated films. (d) The concentration of large clusters in irradiated films both before and after
irradiation. (e) Comparison of large cluster concentration before and after 0.06 dpa irradiation
This work clearly demonstrates the impact Cr has on the incipient stages of defect formation in irradiated Fe. It is also critical to investigate Cr clustering and the presence of ’ precipitation and thus examine its suggested role in enhancing or decreasing radiation resistance in Fe-Cr alloys. To address this point, we carried out APT on Fe-18Cr pristine and irradiated samples.
Atomic Probe Tomography (APT) Analysis
Figure 4.5 depicts visuals of 3D atom maps for the pristine and irradiated Fe-18 Cr specimens. The results revealed a non-uniform distribution of Cr in the pristine Fe-18 Cr films
(Fig. 4.5a), with the composition varying from ~22 % at surface to 12 % at bulk, a variation most likely associated with the grain boundaries. No significant traces of impurities -- C, N, and O -- were found, with corresponding concentrations < 0.07 at. % each. A small Cr2O3 layer near the original sample surface seems to have formed (consistent with the conclusion made earlier based on the behavior of the low-energy positrons); however, the metal directly underneath was not found to be substantially depleted of Cr. In fact, no evidence for Cr clustering in the underlying metal was found, and there was no indication suggesting the formation of any Cr-rich ’ phase in the pristine sample. Similar to the pristine films, the 0.06 dpa irradiated films also exhibited variability in Cr content, ranging from >24 at. % Cr at the surface to ~ 14 at. % Cr in the bulk of the alloy. The presence of a thin surface oxide atop the metal surface was again detected; however, no indication of Cr clustering nor the formation of Cr-rich ’ phase was observed in the irradiated film. 89
Figure 4.5: APT measurements for Fe-18 Cr films. The results are represented by
element-specific 3-D atom maps for (a) pristine films and (b) irradiated films. The elemental
reconstruction maps on the left show projection of atom position in 3D space. On the top right
corner, a 1-D concentration profile created using the composition proximity histogram method
from the 15 at. % Cr iso-concentration surface is also provided for the pristine film. The concentration profile features elemental concentrations as a function of distance from the Fe-Cr
matrix showing the non-uniform distribution of Cr. No traces of other impurities are present in
the films.
As mentioned before, the segregation of Cr and clustering of Cr atoms in Fe-Cr alloys is highly controversial, and their occurrence has been continuously debated, particularly the role of 90 essential parameters like Cr concentration, irradiation temperature, and the nature of the irradiation itself; Cr precipitation is believed to be induced by neutron irradiation rather than ion- irradiation. This is clearly the case here, despite alloying with high Cr composition of 18 wt.%, as confirmed by the APT measurements in Fig. 4.5. However, the incorporation of Cr led to the stabilization of small vacancy clusters to be distinctly detectable by positrons which were not observed for the non-alloyed film (Fig. 4.3). Previous theoretical studies featuring ab-initio calculations and molecular dynamics have discussed the interaction of Cr with point defects [29,
32-35]. These studies have shown that small-sized vacancy defects like monovacancy and divacancies can attract several Cr atoms around them in energetically favored configurations
[36]. Some modeling studies also revealed attractive binding between Cr solutes and vacancies
[37-40]. In the literature, there has also been mention of a solute-drag mechanism to explain radiation-induced segregation, focused mainly on minor impurity elements like P, S, and B where they form tightly bound solute-defect complexes and diffuse significantly in the material before dissociation; however, this has not really been explored experimentally for major alloy components like Cr [41, 42]. All these modeling efforts and predictions in the literature corroborate our experimental observations of vacancy clusters being stabilized in the presence of
Cr.
It is important to recall that our original pristine films were highly porous and defective and that some of the behavior we describe may depend on that original microstructure. While it is hard to comment on whether the addition of Cr would actually facilitate the formation of new small vacancy clusters in otherwise defect-free samples, however, it is clear from our measurements that if small-sized vacancy defects are present, the addition of Cr stabilizes them enough so they become available as sinks for radiation-induced defects Our measurements 91 clearly suggests that the addition Cr stabilizes the presence of vacancy clusters. By combining results from both PAS and APT, we show that Cr does not provide direct sinks for interstitials; rather, the incorporation of Cr stabilizes the presence of vacancy clusters formed in their vicinity by binding with them and forming Cr-vacancy defect complexes which then act as sinks for radiation-induced vacancies and interstitials.
Conclusions
The effect of Cr alloying on the formation of atomic-scale vacancy clusters in Fe was investigated by depth-resolved PALS and DBS in pristine and ion-irradiated thin films while the presence of Cr clustering and the formation of ’ was investigated by APT. The level of Cr was varied from zero to 8% to 18% in an attempt to understand the reason behind the distinct effect of low Cr percentage (less than 10%) and high percentage Cr (higher than 10%) in material response to irradiation. The main difference that the varying Cr content has on the vacancy structure of the material that is revealed from these measurements is that increasing Cr concentration leads to non-uniform distributions in the size and density of vacancy clusters across the film, disrupting the ability of these clusters to act as defect sinks. Further, ’ does not form even at high Cr concentration and thus is not responsible for this distinct effect or for the mechanism of interaction of Cr atoms with radiation-induced defects. This study also reveals that
Cr atoms do not provide a direct sink for interstitials, but they form defect complexes involving both Cr atoms and vacancy clusters, which then act as sinks for radiation-induced defects.
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CHAPTER V. DEFECT STUDIES IN CE:YAG
Introduction
Yttrium Aluminum Garnet, Y3Al5O12 (YAG), is a widely investigated material because of its applications as detectors, laser host materials, and phosphors [1-7]. YAG exists as a complex cubic structure oxide in which Al3+ ions occupy tetrahedral and octahedral sites in the ratio of 3:2, whereas Y3+ ions occupy dodecahedral sites. When YAG is doped with rare-earth ions like Nd3+, Ce3+, Yb3+ etc., they replace Y3+ ions at these dodecahedral sites [8, 9] since their ionic radii values are close.
Figure 5.1: a) Unit cell of YAG crystal. Blue, grey, and red sites are occupied by yttrium,
aluminum, and oxygen, respectively. The dataset is taken from the online public library:
www.materialsproject.org 99
A great deal of research has been conducted on doping and co-doping of YAG with transition metals like Ce, Nd, Eu, Dy, Yb, Cr and investigation of their optical properties [1-16].
Cerium-doped YAG (Ce: YAG) has received considerable attention because of its intense emission at 525 nm and its essential role in InGaN light-emitting diode to convert its blue emission to white light [17-20]. It absorbs part of the blue light from InGaN and emits yellow light, which combines and appears white. Photoluminescence (PL) spectroscopy is the most conventional and reliable method among the scientific community to measure visible light emission when photo excited. What makes it unique is its contactless and non- destructive probing of the electronic structure of materials. PL characteristics of Ce: YAG are very well- investigated. However, the information on the temperature dependence of photoluminescence is limited to a few studies focused only on a single crystal, which reports the conventional trend that PL emission decreases with the increase in temperature [21-24]. However, to the best of the author’s knowledge, there is no previous report in the literature investigating the effect of temperature on the photoluminescence properties of Ce: YAG nanophosphors and transparent ceramics.
Ce: YAG nanophosphors (NPs) offer advantages like small size, less Ce segregation, and simple preparation and can prove to be extremely useful for white light-emitting diodes
(WLEDs) [25-29]. The traditional solid-state chemical methods used to synthesize conventional
Ce: YAG ceramic powder require very high temperatures and prolonged mixing time, often resulting in inhomogeneity and a large grain size of hundreds of microns [30, 31]. Solid-state methods may also induce a high concentration of defects which affect the optical properties and exciton dynamics. Sol-gel methods using wet chemical routes have significantly reduced the 100 temperature required to form a pure garnet phase in YAG. They also reduced the mixing temperature, allowing better control for particle size and doping limit [26, 30-34].
A great deal of research is being done on Ce: YAG to shape the WLED’s technology by carefully tailoring the characteristics of defect centers and electron traps. However, most of the defects studies on Ce: YAG in literature have been primarily focused on single crystals [35-37].
There are only a few studies on the defects in Ce: YAG transparent ceramics which reports the presence of various kind of lattice defects including vacancy clusters and grain boundaries [38].
Careful tailoring of these defects may prove to help tune the luminescence properties of these materials. Still, much work needs to be done to understand better the defects and their influence on the optical properties of these systems. Thermoluminescence (TL), also known as thermally stimulated luminescence (TSE), is one such very effective method that can unearth the presence of traps and also their depths in materials [39, 40].
The first part of this chapter addresses the physical and optical properties of Ce: YAG
NPs prepared by the sol-gel method, including particle size and photoluminescence and the effect of annealing temperatures and environment on them. The results presented here reveal how the luminesce properties of Ce: YAG NPs are strongly affected by annealing temperature and atmosphere. Then it compares their emission with Ce: YAG transparent ceramics and single crystals and investigates their luminescence decay. The temperature dependence photoluminescence (PL) has not been well investigated, and the thermal quenching characteristics of luminescence in Ce: YAG are not yet understood despite its importance in LED performance. This chapter also presents a comprehensive study for temperature-dependent PL from 83 to 633 K for Ce: YAG NPs, single crystals (SCs) and transparent ceramics (TCs). A significant difference between their thermal quenching characteristics is observed and ascribed 101 due to the presence of a large population of trap levels in the band gap and the contribution of thermo-luminescence emission in PL spectra. In contrast to the observed luminescence quenching in NP and SC, the measurements reveal a significant increase in PL with temperature in transparent ceramics.
The second part of this chapter investigates a novel, unusual luminescence we recently observed in Ce: YAG TCs. Five different microstructures of one material, cerium doped yttrium aluminum oxide garnet (Ce:YAG), study their temperature-dependent photo-luminescence (TD-
PL) kinetics using scanning electron microscopy (SEM), positron annihilation lifetime spectroscopy (PALS), thermally stimulated luminescence (TSE), and temperature-dependent photoluminescence (TD-PL) measurements. It is concluded that the presence of large open volumes or vacancy clusters in the 10m grain size TC led to the formation of traps widely distributed in energy levels providing a high density of both shallow and deep traps in the bandgap.
Experimental Methods
Sample Preparation
Two sets of 0.5% Ce doped YAG transparent ceramics were used in this work, one obtained from Baikowski Inc., USA (3m grain TC), and the other (10m grain TC) was grown by vacuum sintering method at our collaborator at Jiangsu key laboratory, China [41]. The mixtures were dried, calcined, and dry-pressed into 25 mm diameter pellets and cold iso- statically pressed under a pressure of 250 MPa. The pellets were then sintered at 1650 °C in flowing oxygen gas for 2 hours. After sintering, the sample was hot isostatic pressed (HIP) at
1550 °C for 2 hours under a pressure of 200 MPa in argon. Then the sample was annealed in air at 1200oC for 20 hours. Ce: YAG nanophosphors were prepared using the sol-gel method using 102 yttrium nitrate, aluminum nitrate, and cerium nitrate mixed in a stoichiometric amount (5% Ce- doped) and dissolved in distilled water. Urea and poly vinyl alcohol were used as complexing and polymerizing agents, respectively. The solution was heated for 2 h at 150°C, dried, and calcined at 600°C for a few hours. Annealing was done in a GSL-1700X series tube furnace in the air at various temperatures from 600 to 1500°C, each for 18 hours. The single crystal (SC) was purchased from Crytur Inc. and is fabricated by the Cz method. The Ce:YAG bulk- phosphors were purchased from Intematix with Ce concentration around 2% and average grain size around 3µm.
X-ray Diffraction (XRD) studies
Structural characterization of the Ce:YAG samples was carried out via X-ray diffraction
(XRD) measurements. The diffraction intensity vs 2θ was recorded using a Rigaku diffractometer, and the average grain size was calculated using the Scherer equation. The observed patterns were then analyzed using MDI JADE software.
Scanning Electron Microscopy (SEM)
An FEI Quanta 3D FEG scanning electron micrograph (SEM) system capable of imaging in high vacuum, low vacuum and ESEM modes were used to obtain SEM micrographs of Ce: YAG transparent ceramics. A thin layer (approx. 15 nm) of gold was sputter-coated on the samples to make them conductive for the SEM analysis.
Photoluminescence (PL) Spectroscopy
PL measurements were carried out using a 456 nm excitation band delivered from an
LED connected to a monochromator to remove any tail of longer wavelength that may be emitted from the LED. The PL emission was recorded by a CCD sony detector and diffraction grating covered from 200 to 800 nm. Temperature-dependent photoluminescence studies were 103 carried out in the range of 83- 640 K by recording PL emission versus wavelength at every 20-30
K interval using an INSTECH Heat Stage Profile instrument. The emission intensity was then integrated and plotted versus the corresponding temperature.
Thermal Stimulated Emission (TSE) Spectroscopy
Thermal stimulated emission (TSE) measurements were carried out via a special spectrometer designed and constructed in-house, which allows emission to record as a function of both wavelength and temperature [42-44]. To examine the contribution of TSE in temperature-dependent PL emission, the same 456 nm blue LED was used for irradiation at -
1900C for 30 minutes and the samples were linear ramp heated to 640 K at a rate of 60 °C/min.
For the representation of complete trap characteristics, a Xe lamp was used as a TSE irradiation source. Shallow trap measurements were carried out by irradiating the samples at -1900C for 30 minutes, which were then linear ramp heated to room temperature at 60 °C/min. Deep traps were identified by irradiation of samples at room temperature, which were then linear heated to 640 K at a rate of 60 °C/min.
Positron Annihilation Spectroscopy (PAS)
Positron annihilation lifetime measurements (PALS) were performed using gamma- induced positron spectroscopy (GIPS) at the ELBE (Electron Linac with high Brilliance and low
Emittance) facility, at the Helmholtz-Zentrum Dresden-Rossendorf (HZDR) in Dresden,
Germany [45-47]. GIPS is an advanced PAS technique that can generate a positron decay curve free from background or source contributions. It uses high-energy γ-rays to generate positrons directly inside the sample by pair production. The main advantage is that it completely eliminates unwanted contributions from positron annihilation in either the source or cladding materials and thus results in accurate measurements of positron lifetimes. The positron lifetime 104 decay curves of Ce: YAG TCs and single-crystal are generated. GIPS at ELBE facility uses a superconducting electron accelerator in the continuous wave (CW) mode with a high repetition rate of 26 MHz, resulting in intense high-energy -rays that produce a sufficient number of positrons inside the sample through pair production. The timing resolution of the accelerator is about 5ps, resulting in very short pulses and thus providing an excellent opportunity to perform
PALS with good timing resolution in most materials. The measurements were carried out by a coincident measurement - Age Momentum Correlation (AMOC) - of the time of arrival and energy of the annihilation photons. This dramatically reduces the background of scattered photons resulting in spectra with an excellent signal-to-background ratio. The positron lifetime is registered as the time difference between the creation of positron in the sample (indicated by the accelerator pulse) and annihilation photons. In the current measurements, the electron-beam parameters were pulse width 5 ps, current 600 μA, accelerator energy 16.9 MeV, repetition rate
26 MHz, and a 12.5 μm thick Nb foil was used as a radiator to create the bremsstrahlung beam.
Two AMOC detection systems were employed to collect the annihilation photons from the sample simultaneously to ensure the consistency of the results.
Results And Discussions
Characterization Of Ce:YAG NPs
Figure-5.2 shows the X-ray diffraction pattern of Ce: YAG NPs annealed at various temperatures ranging from 1000 to 1500°C. The formation of the pure YAG phase is first seen at
1000°C. However, when the sample was annealed in the range of 600 to 900°C, the XRD pattern revealed the absence of the YAG phase. When the sample was annealed at higher temperatures of 1000- 1500°C, a pure YAG phase was obtained; however, a minor peak indicating the
o presence of CeO2 impurity phase at 2휃 = 28 was observed. The formation of the CeO2 phase can 105 either be due to phase segregation or non-uniform mixing of the reactants [48]. Figure-5.3 illustrates the dependence of grain size on the annealing temperature. At higher temperatures, the particle size tends to increase with annealing temperature due to grain growth. However, it started to saturate at about 1500°C.
Figure 5.1: XRD patterns of Ce: YAG nanophosphors annealed in various temperatures from 1000 to 1500°C. The un-doped YAG phase (shown in black) was included in the graph to
facilitate comparison. 106
Figure 5.2: Grain size of Ce: YAG nanophosphors annealed at various temperatures from 1000
to 1500oC.
Figure-5.4 displays the photoluminescence (PL) emission spectrum of Ce: YAG NPs.
The well-known emission at 525 nm results from 5d to 4f transitions. No PL was observed after annealing between 600 and 900°C, consistent with the absence of the YAG phase in the XRD patterns. It can be seen that PL intensity dramatically increases with annealing temperature.
There are a few factors that may be responsible for this increase: a) better incorporation of Ce3+ in the matrix, b) more efficient formation of YAG phase, and c) decrease in surface defects possibly resulting from a lower surface area to volume ratio; as the particle size increases, the grain density and the subsequent number of surface defects decrease, however lower annealing temperatures lead to smaller particle sizes, which in turn lead to a higher number of surface defects that may inhibit the transport of photons. It is interesting to note that if we anneal at high temperatures in the presence of air, the process of oxidation is inevitable and some of the Ce3+ ions oxidize to Ce4+. To verify this, Ce: YAG nano-phosphor was annealed in air at 1200°C for 107
12 hours and then divided into three parts. One was left as it is, while the other two parts were separately re-annealed in relatively reduced atmospheres of 1) hydrogen and 2) argon, each at
1200°C for 1 hour. Their PL spectra in Fig. 5.4b show that the PL intensity significantly increased after annealing in reducing atmospheres of H2 or Ar. The reason behind this is the reduction of Ce4+ to Ce3+ state, which is responsible for the yellow emission in YAG.
Figure 5.4: PL intensity of Ce: YAG nanophosphors as a function of wavelength. (a) for different
annealing temperatures, (b) for different annealing atmosphere
Figure-5.5 compares the PL emission of Ce:YAG NPs (annealed at 1500°C) with Ce:
YAG single crystal (SC) and transparent ceramics (TC) emission recorded under the same excitation conditions. It is interesting to note that the Ce: YAG NPs prepared in this work are more efficient than SC and transparent ceramics despite the excellent transparency of the Ce:
YAG transparent ceramic samples. No significant peak shift is observed in nano-phosphor
(λmax=533nm) compared to ceramics (λmax=532nm) and single crystal (λmax=527nm). The high intensity of PL emission in Ce: YAG NP can be explained due to the high doping of Ce (5%) compared to (0.1%) in single crystal and (0.5%) in transparent ceramics. Both SC and transparent ceramics have been known to face strong Ce segregation and agglomeration when 108 doped with high concentrations. One of the advantages of the wet chemical method for phosphor preparation is that it allows higher doping. Figure-5.5 demonstrates the high efficiency of the Ce:
YAG NPs prepared in this work and its relevance to be used in the investigation of the luminescence characteristics of NPs.
Figure 5.5: PL emission intensity as a function of wavelength. The graph compares the peak
emission and shape of Ce: YAG NPs with single crystals and transparent ceramics.
A main drawback of phosphors is their luminescence quenching with a minor increase in temperature, limiting their applications. Though many groups have tried to address the issue in bulk materials [48, 49], much has to be done to utilize nanophosphors as a substitute emitter with excellent luminescence even at high temperatures. Up to now, there is no significant data on luminescence quenching with increasing temperature in Ce: YAG nanophosphors. To address that and to obtain a better understanding of PL characteristics in Ce: YAG, PL spectra of Ce:
YAG NPs (annealed at 1300°C) were recorded at various temperatures from 83 to 623 K (Fig. 109
5.6a). The spectra demonstrate intense emission at 83 K, which gradually decreases with increasing temperature. However, Fig. 5.6b shows that even at a high temperature of 623 K, Ce:
YAG NPs prepared in this work still emit low light levels. At low temperatures, the emission band displays the well-known double band structure peaks (Fig. 5.6a), which broaden and appear as one peak at higher temperatures [48]. These double band peaks result from transitions from the lowest Stark level of the 5d excited state to the two Stark levels (2 F5/2 and 2F7/2) of the 4f ground state of Ce3+.
The luminescent decay with increasing temperature is often observed in materials, but an exact universal explanation has yet to be reported. It can be, however, related to the temperature dependence of the oscillator strength of a material. In Ce: YAG, it is expected that 2 F5/2 levels become thermally populated at higher temperatures, resulting in less absorption strength and less luminescence output. Also, another possible reason for the decrease in luminescence is the energy migration to defects and non-radiative decay. PL emission spectra were recorded from 83 to 623 K at every 20 or 30 K interval, Fig. 5.6a includes only a handful of them for better display, while Fig. 5.6c displays the integrated peak area as a function of temperature measured from 83 to 623 K to give the overall picture. The graph reveals an exciting behavior. Instead of gradually quenching with the increase in temperature, PL intensity undergoes a slight increase around 200 K, then decreases again.
To provide a deeper insight into the characteristics of PL emission intensity decrease with temperature, temperature-dependent photoluminescence (TD-PL) measurements were carried out on Ce: YAG SC and transparent ceramics, followed by TSE spectroscopy. 110
Figure 5.6: TD-PL kinetics of NPs. (a) TD-PL emission measurement of NP as a function of wavelength over a temperature range of 83K-623K. (b) PL emission intensity of NP at T=623K.
(c) Integrated PL peak intensity as a function of temperature for NP. It shows that the luminescence intensity sharply decays with rising temperature up to 200 K, where it undergoes a
little increase, then decays again.
Figures 5.7 and 5.8 illustrate the temperature dependence of PL in SCs and transparent ceramics, respectively. Single crystals and NPs share the same characteristics, including the slight increase at 200 K, and exhibit overall decay in luminescence with rising temperature while
PL in Ce: YAG transparent ceramics manifest surprising behavior in which the luminescence 111 intensity is not maximized at low temperatures; on the contrary, it becomes very intense and broad in the 373- 473 K temperature range before quenching.
Figure 5.7: TD-PL kinetics of SC. a) TD- PL measurements of Ce: YAG single crystal. b)
Integrated PL intensity of Ce:YAG single crystal as a function of temperature.
Figure 5.8: TD-PL kinetics of TC. a) Temperature-dependent PL emission spectra of Ce:
YAG transparent ceramics as a function of wavelength over a temperature range of 83 to 633 K.
b) Integrated PL peak intensity of Ce: YAG transparent ceramics as a function of temperature.
Previous studies on TSE have proved it to be an effective tool to investigate trap levels in semiconductors and dielectrics [39, 40] and correlate them with many of the electronic properties of the materials. TSE behavior in Ce: YAG may facilitate interpretation of the TD-PL behavior 112 reported in this work. The interest in the current study is to investigate the possible presence of
TSE emission and its contribution to the measured PL intensity. The samples were illuminated with 456 nm light, the same wavelength we used for PL excitation to populate electrons in the conduction band, at a temperature of 83K. No TSE emission was detected in Ce: YAG NPs or single crystals, while a strong TSE emission was measured in Ce: YAG transparent ceramics.
Figures 5.9a and 5.9b display the TSE contour plot and glow curve, respectively. The glow curve was constructed by integrating the emission over all wavelengths at each temperature. By comparing the TSE glow curve in Fig. 5.9b with the PL emission in Fig. 5.8b, it can be concluded that TSE emission significantly contributes to the PL intensity and is responsible for the increase of PL with temperature. This conclusion is also consistent with the absence of TSE emission in NPs and SCs, and their PL emission intensity decreases.
Figure 5.9: TSE emission in Ce:YAG TC. a) Thermo-luminescence contour plot of Ce: YAG
transparent ceramics as a function of temperature and wavelength. Excitation was carried out
using a sub-band gap light of 456 nm (2.72 eV). b) Glow curve of TSE emission of Ce: YAG
transparent ceramics. 113
The 456 nm (2.72 eV) light used for excitation in TSE and PL is much lower than the band gap of YAG, which is in the order of 6.5 eV and thus cannot excite electrons from the valence band to the conduction band. A defect center below the conduction band by 2.72 eV should provide electrons to the conduction band, while a separate defect center or impurity may act as an electron trap. From the analysis of the TSE glow curve, we identified a trap level of
1.021 ± 0.074 eV below the conduction band. From these calculations and by comparing emission in fig. 8b and 9b, we suggest a model (Figure-10) explain the exciting behavior of PL increase with temperature. The model illustrates a four-step mechanism: a) excitation of an electron from a defect center to the conduction band, b) capture of the electron by a trap level of
1.02 eV in the band gap, c) thermal stimulation of the electrons from the trap, d) electron recombination in Ce atom emitting luminescence. No TSE emission has been detected in NPs and SCs after excitation by 456 nm light. Accordingly, the slight increase in PL around 200 K cannot be explained by this model; it is most likely a result of thermal excitations of electrons from 4f to 5d levels during heating which then decays back and leads to a modest increase in PL intensity despite the contrary effect of the temperature dependence of the oscillator strength. We believe that the big difference between NP/SC and transparent ceramics is their microstructure.
The presence of voids in the grain boundaries of transparent ceramics provides an array of trap levels leading to this behavior. Previous positron annihilation spectroscopy measurements on
YAG and Ce: YAG [21, 48] have confirmed the presence of large voids only in transparent ceramics samples. 114
Figure 5.10: Diagram illustrating the mechanism behind TL emission and its contribution to PL
spectrum. a) excitation of an electron from the defect center to the conduction band by 456 nm
light. b) trapping of the electron by another defect center. c) thermal stimulation. d)
recombination of charge carriers and light emission.
Investigation Of TD-PL Kinetics and Trap-Assisted Luminescence
Figure 5.11 depicts the TD-PL kinetics of five Ce:YAG samples with different microstructures (single crystal, transparent ceramic with 10m average grain size, transparent ceramic with 3m average grain size, nano-phosphor with 8nm average grain size, bulk phosphor with 10m average grain size). The TD-PL intensities are normalized to the PL intensity at the first measurement point at 83 K. The TD-PL for the NP and single-crystal samples exhibit the traditional behavior of luminescence quenching with increasing temperature.
However, the decay rate is very different, as shown in Fig. 5.11a and 5.11b. In contrast, a considerable rise in PL intensity with increasing temperature was recorded from the 3m grain
TC, 10m grain TC and bulk phosphor, as shown in Figs. 5.11c, 5.11d, and 5.11e, respectively.
With increasing temperature, the PL emission increases around 250% for the 3m grain TC and
20 to 25% for the 10m grain TC and bulk phosphor. The stark difference in PL kinetics 115 observed for these different systems suggests that this behavior is not related to the Ce atom or
YAG structure but is a characteristic of the system itself and the defect structure associated with it. The increase in emission with temperature in the garnet ceramics and the bulk phosphor is opposite to the well-known decay of PL, and it contradicts the temperature dependence of PL in
Ce: YAG single crystals and nanophosphors. The PL from the TCs features a plateau from 90 to
400 K in the 10m grain TC and 300-500 K in the 3m grain TC. The apparent difference between these five systems for the same material is their microstructures. We anticipate that this difference may lead to traps with different spatial and energy distributions. These traps store energy (trap charge carriers) during the PL excitation through certain mechanism(s) and release them later to the Ce atom gradually with increasing temperature over imposing on the PL emission and modifying the PL temperature dependence of PL kinetics with temperature.
Figure 5.11: TD-PL kinetics in various Ce:YAG microstructures. Temperature-dependent
photoluminescence measurements of Ce:YAG samples in the range 83- 633K for (a) Nano-
phosphor, (b) Single-crystal, (c) Bulk-phosphor, (d) 3m avg grain size TC and (e) 10m avg 116 grain size TC. The TD-PL intensities are normalized to the PL intensity at the first measurement
point at 83 K. The excitation wavelength used was 456 nm.
To investigate this notion and reveal the mechanism behind this phenomenon, thermally stimulated emission (TSE) was applied, which can reveal the presence of traps and their distributions. First, the TSE measurements were performed using 456 nm light (the same wavelength used for excitation in the PL experiment) for 30 mins at 83 K. The emission was recorded as a function of wavelength and temperature (in the same temperature range) using an in- house spectrometer. TSE glow curves were generated by integrating emission intensity over all wavelength range (200- 670 nm) at each temperature and presented in Fig. 5.12. The figure shows an emission peaking around 500°C arising from a single trap level in the 3m grain TC and multiple emissions from the 10m grain TC, indicating contributions from several trap levels. As mentioned above, the TSE glow curves in Fig. 5.12 have been generated using one excitation wavelength of 456 nm, which only excites carriers from energy levels of 2.72 eV or less from the bottom of the conduction band (but not from the valance band) and prompts them to the conduction band where they may be trapped at defects, then released later upon heating transferring their energy to Ce ion and producing emission. This suggests that this process may happen similarly in temperature-dependent PL measurements and contribute to PL emission later at a higher temperature and modifies the kinetics of PL with respect to the temperature.
117
Figure 5.12: Comparison of thermoluminescence glow curves of 3m and 10m avg grain size
TCs. Irradiation for TL was provided using 456 nm blue LED light, the same LED used for
excitation in temperature-dependent PL measurements.
Both Ce:YAG SC and NP showed no TSE emission when excited with 456 nm blue LED light which is expected based on their normal PL decay curves with temperature. Surprisingly, the Ce:YAG BP did not reveal any TSE emission when illuminated with 456 nm despite their unusual temperature-dependent kinetics.
To search for a mechanism responsible for the rise of PL with temperature in Ce:YAG
TC and Ce:YAG BP, the presence of traps and their distributions and characteristics are investigated. This can be done through TSE measurements using a high-energy excitation source to pump electrons from the valence band to the conduction band. Fig. 5.13 presents the glow 118 curves from 83 to 650 K for all samples. The samples were excited at 83K using a Xenon lamp which broadband light extending to deep ultraviolet providing high energy excitation above 6.5 eV, the band-gap energy of YAG. Once electrons are pumped to the conduction band, they may fill the traps. However, they can be released gradually to the Ce atom by thermal stimulation.
The TSE glow curves in Fig. 5.13 represent a complete representation of the trap levels and characteristics in each sample. The symbols represent the experimental data, and solid lines represent the fit. The Ce: YAG NP did not give any TSE emission probably because of the large light scattering at the grain boundaries, while Ce:YAG SC and Ce:YAG BP exhibited intense
TSE in the low-temperature range. However, only weak emission was observed at high temperatures for both Ce:YAG SC and Ce:YAG BP.
In contrast, the glow curves for both Ce:YAG TCs, 3m grain TC, and 10m grain TC exhibited TSE peaks throughout the temperature range, indicating the presence of both shallow trap levels and a series of deep traps. This may be due to the presence of a large number of grain boundaries in the Ce:YAG TCs. Nonetheless, the nature and shape of the glow curves obtained for both TCs were quite distinct as commercial TC showed several relatively sharp peaks at low temperatures compared to a single broad peak observed for 10m grain TC. 3m grain TC displayed an intense broad peak spanning from 400- 600 K compared to two broad peaks for
10m grain TC at high temperatures. This suggests that both TCs possess a series of trap levels with different origins and characteristics, which plays a vital role in upholding the luminescence at a similar level for a broad span of temperature even after reaching its emission maximum. 119
Figure 5.12: TSE emission in various Ce:YAG microstructures. TSE glow curves depicting
intensity versus temperature in temperature range from 80 to 640 K representing both shallow
and deep traps for (a) Single-crystal, (b) Bulk-phosphor, (c) 3m avg grain size TC, (d) 10m
avg grain size TC. The experimental data points are represented as symbols, while solid lines
demonstrate the gaussian fit of the glow curves. The illumination source was a broad-band Xe
lamp.
The obvious difference between the glow curves of SC, which features the traditional PL decay with temperature, and the TCs and bulk phosphor, which feature the rise of PL emission with temperature, is the characteristics of glow peaks. The Ce:YAG SC exhibits narrow individual peaks (4 peaks in the low-temperature range) corresponding to four shallow traps associated with 120 lattice defects in the lattice and two weak peaks in the high-temperature range corresponding to deep traps. However, the glow curves of all TCs and bulk phosphors exhibit very broad peaks indicating the presence of series of a large number of traps distributed throughout the bandgap.
This series formed of a large number of traps closely distributed in energy is responsible for the reversal behavior of PL with temperature observed in TCs and bulk phosphors.
To understand the origin behind the formation of such a high density of traps with different energy levels and a deeper insight into the nature of traps, positron annihilation spectroscopy
(PAS) was employed. Fig. 5.14 shows the lifetime spectra measured by GIPS for the two Ce:
YAG TCs and single-crystal indicating different positron decay curves. Their corresponding average positron lifetime values are tabulated in Table 5.1. The average positron lifetime values obtained from the spectra are 207 + 1 ps for the 10m grain TC and 190 + 1 ps for the 3m grain
TC compared to 168 + 1 ps for the Ce:YAG SC. The large average lifetimes in the Ce:YAG TCs indicate the presence of a large fraction of defect clusters, which is absent in single crystal and bulk phosphor. PALS confirms the presence of large defects clusters in the Ce:YAG TCs, which are most likely the reasons for the trap centers detected by TSE even when illuminated by 456 nm light. Thus, it is possible that the nature and distribution of these clusters strongly affect exciton dynamics and contribute to PL emission at higher temperatures. 121
Figure 5.14: GIPS measurements on Single-crystal, 3m avg grain size TC and 10m avg grain size TC. (a) Positron lifetime decay spectra (b) Table showing corresponding values of average
positron lifetimes. A larger lifetime value indicates the presence of large defect clusters.
Table 5.1: Table showing corresponding values of average positron lifetimes. Larger lifetime
value indicates the presence of large defect clusters.
Samples Average e+ lifetimes (ps)
Single crystal 168 ± 1
3m grain sized TC 190 ± 1
10m grain sized TC 207 ± 1
122
Conclusions
The physical and luminescence properties of Ce:YAG NPs and their dependence on annealing temperature and atmosphere were investigated. The measurements showed that
Ce:YAG NPs prepared by simple sol-gel methods are highly efficient phosphors. High- temperature annealing was found to increase the grain size and enhance the green luminescence, while annealing in reducing atmospheres of Ar or H2 significantly improved the emission. TD-
PL measurements revealed an exciting luminescence phenomenon in transparent ceramics, which leads to a significant increase in luminescence with temperature. This phenomenon has not been previously reported. The presence of voids in transparent ceramic generates an extensive array of trap levels triggering this behavior; thus, we suggest that this new luminescence phenomenon is characteristic of transparent ceramics.
In addition, this work further explored the interesting and surprising TD-PL kinetics in various microstructures of Ce:YAG systems. It illuminates the novel trap-assisted luminescence aided by defect clusters in transparent ceramics, facilitating an increase in PL intensity at higher temperatures. The absence of this phenomenon in Ce: YAG single crystals and nanophosphors indicate that it is not associated with Ce emission or garnet structure, but it is unique to garnet transparent ceramics. It will be interesting to study other transparent ceramics to examine if this phenomenon extends beyond garnet and if it is characteristic of transparent ceramics in general.
Lastly, increased luminescence with temperature in garnet TCs reported in this work may open up new interesting applications.
123
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CHAPTER VI. SUMMARY AND CONCLUSIONS
The studies reported here in this thesis focus on the powerful capabilities of PAS in investigating point defects at the atomic-level in different material systems: metal (Fe), alloys
(Fe-8Cr and Fe-18Cr), and oxide (Ce:YAG) and demonstrate its effectiveness for developing materials for a wide range of applications, such as in nuclear reactors, radiation sciences and photonics respectively.
By employing depth-resolved PAS in Fe, we tried to enhance the understanding of irradiation- induced defects at the atomic-scale. The irradiation was carried out at different doses in the range
0-0.6 dpa, and the evolution of defects was studied quantitatively using PAS and TEM. The measurements showed that the increase in the density of small vacancy clusters with irradiation is associated with a remarkable reduction in the size of large voids, revealing a novel mechanism for the interaction of cascade damage with voids in ion-irradiated materials, a consequence of the high porosity of the initial microstructure.
In addition, by combining depth-resolved PAS and APT techniques, we further investigated the effect of Cr alloying on the formation and evolution of atomic size clusters induced by ion irradiation in Fe. The study reveals that the well-known resistance to radiation in
Fe-Cr alloys arises from the stabilization of vacancy clusters around Cr atoms which act as sinks for radiation-induced defects. Thus, Cr atoms do not provide a direct sink for interstitials; rather, they form defect complexes of Cr atoms and vacancies, which then act as sinks for irradiation- induced vacancies and interstitials. Most importantly, we find that lower amounts of Cr create smaller, uniformly distributed defect clusters that act as efficient sinks for radiation damage, but larger quantities of Cr form a defect structure that is less homogenous in size and spatially distribution resulting in less efficient damage recombination. No evidence of phase ’ was found 130 before or after irradiation, which indicates that it does not play a role in enhancing or hindering radiation tolerance.
Lastly, this work reports an interesting temperature-dependent photoluminescence (TD-
PL) behavior in an important photonic oxide material, Ce:YAG, where luminescence increases with the increase in temperature as opposed to traditional luminescence quenching processes.
Here, we apply PALS to study defect characteristics in several different microstructures of
Ce:YAG, and we show that TD-PL kinetics can be controlled by modifying microstructure and engineering defects. Combining PL measurements with TSE, SEM, and PAS, we conclude the presence of large open volumes or vacancy clusters in Ce:YAG transparent ceramics, which results in trap levels widely distributed in the bandgap.
131
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156
APPENDIX A. PUBLICATIONS
1. Hernandez, A., Islam, M.M., Saddatkia, P., Codding, C., Dulal, P., Agarwal, S., Janover,
A., Novak, S., Huang, M., Dang, T. and Snure, M. MOCVD growth and characterization
of conductive homoepitaxial Si-doped Ga2O3. Results in Physics, 2021, 25, 104167.
2. Agarwal, S., Liedke, M. O., Jones, A. C. L., Reed, E., Uberuaga, B. P., Wang, Y.Q.,
Cooper, J., Kaomi, N., Li., N., Auguste, R., Hosemann, P., Capolungo, L., Edwards, DJ,
Butterling, M., Hirschmann, E., Wagner, A., Selim, F. A. A new mechanism for void-
cascade interaction from nondestructive depth-resolved atomic-scale measurements of
ion irradiation-induced defects in Fe. Sciences Advances. 2020, 6, eaba8437.
3. Saadatkia, P., Agarwal, S., Hernandez, A., Reed, E., Brackenbury, I.D., Codding, C.L.,
Liedke, M.O., Butterling, M., Wagner, A. and Selim, F.A. Point and extended defects in
heteroepitaxial β− G a 2 O 3 films. Physical Review Materials, 2020, 4, 104602.
4. Islam, M.M., Adhikari, N., Hernandez, A., Janover, A., Novak, S., Agarwal, S., Codding,
C.L., Snure, M., Huang, M. and Selim, F.A. Direct measurement of the density and
energy level of compensating acceptors and their impact on the conductivity of n-type
Ga2O3 films. Journal of Applied Physics, 2020, 127, 145701.
5. Zhang, L., Wu, J., Stepanov, P., Haseman, M., Zhou, T., Winarski, D., Saadatkia, P.,
Agarwal, S., Selim, F.A., Yang, H. and Zhang, Q. Defects and solarization in YAG
transparent ceramics. Photonics Research, 2019, 7, 549-557.
6. Agarwal, S., Haseman, M.S., Leedy, K.D., Winarski, D.J., Saadatkia, P., Doyle, E.,
Zhang, L., Dang, T., Vasilyev, V.S. and Selim, F.A. Tuning the Phase and
Microstructural Properties of TiO2 Films Through Pulsed Laser Deposition and 157
Exploring Their Role as Buffer Layers for Conductive Films. Journal of Electronic
Materials, 2018, 47, 2271-2276.
7. Agarwal, S., Haseman, M.S., Khamehchi, A., Saadatkia, P., Winarski, D.J. and Selim,
F.A. Physical and optical properties of Ce: YAG nanophosphors and transparent ceramics
and observation of novel luminescence phenomenon. Optical Materials Express, 2017, 7,
1055-1065.
8. Selim, F.A., Khamehchi, A., Winarski, D. and Agarwal, S. Synthesis and characterization
of Ce: YAG nano-phosphors and ceramics. Optical Materials Express, 2016, 6, 3704-
3715.
158
APPENDIX B. CONFERENCE CONTRIBUTIONS
1. S. Agarwal, M. Haseman, A. Khamehchi, F. A. Selim “Temperature dependent
luminescence characteristics of Ce: YAG nanophosphors and transparent ceramics and
observation of novel phenomenon”, OSAPS/MI-AAPT 2018 Spring meeting, Michigan
State University, East Lansing, Michigan, USA
2. S. Agarwal, L. Zhang, P. Stepanov, M. Haseman, F. A. Selim, “Effect of Microstructures
on Luminescence Kinetics in Transparent Ceramics”, OAS (Ohio Academy of Sciences,
2018) 2018 annual meet, Bowling Green State University, Ohio, USA
3. S. Agarwal, L. Zhang, M. Butterling, M. Liedke, A. Wagner, F. A. Selim, “Effect of
microstructures on luminescence kinetics in Ce: YAG transparent ceramics”, ICPA- 18
(18th International conference on positron annihilation, 2018), Orlando, Florida, USA
4. S. Agarwal, M. Haseman, A. Khamehchi, P. Saadatkia, D. J. Winarski, F. A. Selim,
“Physical and optical properties of Ce:YAG nanophosphors and transparent ceramics and
observation of novel luminescence phenomenon”, Ohio Photochemical Society (OoPS)
Meeting 2017, Oregon, Ohio, USA
5. S. Agarwal, M. Haseman, P. Saadatkia, D. Winarski, E. Doyle, L. Zhang, K. Leedy, F. A.
Selim, “Effects of substrate type and temperature growth on the microstructure and
characteristics of pulsed laser deposited TiO2 films and their role as buffer layers for
conductive films”, EMC- 59 (59th Electronic Materials Conference, 2017), University of
Notre Dame, Indiana, USA
6. S. Agarwal, L. Zhang, M. Haseman, P. Stepanov, F. A. Selim, “Development of positron
annihilation spectroscopy as a powerful tool for studying solarization phenomena and
advancing transparent ceramics technology”, JPos (International Workshop on Physics 159
with Positrons at Jefferson Lab, 2017), Thomas Jefferson National Accelerator Facility,
Virginia, USA
7. S. Agarwal, F. A. Selim “Effect of particle size on nature of defects and luminescence
characteristics in Ce: YAG phosphors”, ICPA- 18 (18th International conference on
positron annihilation, 2018), Orlando, Florida, USA
8. S. Agarwal, L. Zhang, M. Haseman, F. A. Selim “Effect of doping and sintering
temperature on the nature and energy levels of defects in Ce: YAG transparent ceramics”,
ICPA- 18 (18th International conference on positron annihilation, 2018), Orlando,
Florida, USA